ISBN: 0-8247-0074-0This book is printed on acid-free paper. Headquarters Marcel Dekker, Inc. 270 Madison Avenue, New York, NY 10016 tel: 212-696-9000; fax: 212-685-4540 Eastern Hemisphere Distribution Marcel Dekker AG Hutgasse 4, Postfach 812, CH-4001 Basel, Switzerland tel: 41-61-261-8482; fax: 41-61-261-8896 World Wide Web http://www.dekker.com The publisher offers discounts on this book when ordered in bulk quantities. For more information, write to Special Sales/Professional Marketing at the headquarters address above. Cover illustration: Scanning electron micrograph of an oxide/oxide composite exposed to an environment of air and water at 1000°C. Courtesy of C.A. Lewinsohn, Pacific Northwest National Laboratory. Copyright 2001 by Marcel Dekker, Inc. All Rights Reserved. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming, and recording, or by any information storage and retrieval system, without permission in writing from the publisher. Current printing (last digit): 10 9 8 7 6 5 4 3 2 1 PRINTED IN THE UNITED STATES OF AMERICA Preface New materials, composites, and coatings are being developed at a rapid rate, and there has been an increase in the substitution or replacement of one class of material by another. More complex materials are being engineered and used in a new variety of environments. Many materials are used in composite or coated forms to enhance performance. Various combinations of metals, intermetallics, ceramics, and polymers are becoming more common. Composites with discontinuous, dispersed phases within a matrix, and fiber and laminated reinforcements are being developed. Coatings also include combinations similar to those for bulk composite materials. Materials are being pushed to perform in a wider range of environments than ever before. Aqueous and high-temperature environments, which may contain varying amounts of corrosive species, are commonly encountered by advanced materials. In other cases, the effect of environments such as water, solvents, wine, and food thought to be relatively benign must be understood. All these developments have made it difficult to locate information on the effects of environment on the new materials, composites, and coatings. This comprehensive book describes such effects for a broad range of materials and environments, filling the information gap and providing a comprehensive viewpoint for the scientist or engineer interested in applying new materials to existing applications or old materials to new applications. This book would not have been possible without the many hours given by each contributor. Their effort and dedication are greatly appreciated. Also, the assistance of B. H. Wardlow at PNNL in coordinating the manuscripts is greatly appreciated. Russell H. Jones iii Contents Preface Contributors iii vii I. Metallic Alloys 1. Ferrous Alloys (Ferritic and Martensitic) Bruce Craig 1 2. Austenitic Stainless Steels Russell H. Jones, Stephen M. Bruemmer, Mike J. Danielson, and Bruce Craig 31 3. Nickel-Based Alloys for Resistance to Aqueous Corrosion Paul Crook 55 4. Nickel-Based Alloys for Resistance to High-Temperature Corrosion Mark A. Harper and George Y. Lai 5. Corrosion of Copper and Its Alloys Andrew James Brock 75 115 v vi Contents 6. Reactive and Refractory Alloys Te-Lin Yau 151 7. Aluminum Alloys N. J. Henry Holroyd 173 8. Magnesium Alloys Mike J. Danielson 253 II. Intermetallic Alloys 9. Environmental Embrittlement of Nickel-Based and Iron-Based Intermetallics Norman S. Stoloff III. 275 Ceramics 10. Nonoxide Ceramics Nathan S. Jacobson and Elizabeth J. Opila 311 11. Oxide Ceramics F. S. Pettit, G. H. Meier, and J. R. Blache`re 351 IV. Composites 12. Metal Matrix Composites Russell H. Jones 375 13. Ceramic Matrix Composites Russell H. Jones, C. H. Henager, Jr., Charles A. Lewinsohn, and Charles F. Windisch, Jr. 391 14. Issues in Predicting Long-Term Environmental Degradation of Fiber-Reinforced Plastics Aaron Barkatt 419 V. Metallic Glasses 15. Amorphous and Nanocrystalline Alloys Koji Hashimoto 459 Index 501 Contributors Aaron Barkatt Department of Chemistry, The Catholic University of America, Washington, D.C. J. R. Blache`re Materials Science and Engineering Department, University of Pittsburgh, Pittsburgh, Pennsylvania Andrew James Brock Haven, Connecticut Metals Research Laboratories, Olin Corporation, New Stephen M. Bruemmer Pacific Northwest National Laboratory, Richland, Washington Bruce Craig MetCorr, Denver, Colorado Paul Crook diana Engineering and Technology, Haynes International, Kokomo, In- Mike J. Danielson Pacific Northwest National Laboratory, Richland, Washington Mark A. Harper Research and Development, Special Metals Corporation, Huntington, West Virginia vii viii Contributors Koji Hashimoto Tohoku Institute of Technology, Sendai, Japan C. H. Henager, Jr. Pacific Northwest National Laboratory, Richland, Washington N. J. Henry Holroyd erside, California Research and Development, Luxfer Gas Cylinders, Riv- Nathan S. Jacobson Materials Division, NASA Glenn Research Center, Cleveland, Ohio Russell H. Jones Pacific Northwest National Laboratory, Richland, Washington George Y. Lai Consultant, Carmel, Indiana Charles A. Lewinsohn Washington Pacific Northwest National Laboratory, Richland, G. H. Meier Materials Science and Engineering Department, University of Pittsburgh, Pittsburgh, Pennsylvania Elizabeth J. Opila Department of Chemical Engineering, Cleveland State University, Cleveland, Ohio F. S. Pettit Materials Science and Engineering Department, University of Pittsburgh, Pittsburgh, Pennsylvania Norman S. Stoloff Materials Science and Engineering Department, Rensselaer Polytechnic Institute, Troy, New York Charles F. Windisch, Jr. Washington Pacific Northwest National Laboratory, Richland, Te-Lin Yau Te-Lin Yau Consultancy, Albany, Oregon 1 Ferrous Alloys (Ferritic and Martensitic) Bruce Craig MetCorr, Denver, Colorado I. INTRODUCTION This chapter addresses the corrosion behavior of ferrous alloys, specifically ferritic and martensitic irons and steels. The reason for this designation is to distinguish these alloys from the austenitic alloys that will be discussed in a later chapter. However, the use of the terms ‘‘ferritic’’ or ‘‘martensitic’’ is not intended to exclude pearlitic or bainitic microstructures, but is only intended as a convenience. Therefore, the discussion in this chapter addresses all low-alloy ferrous materials and ferritic and martensitic stainless steels. The largest group of ferrous alloys are steels which will be the emphasis in this chapter, however, cast irons, several of which can be quite corrosion resistant, will also be mentioned. There are tens of thousands of different steels in the world; however, they are usually referred to in groups as a function of their chemical composition. Thus, carbon steels (also referred to as mild steels) contain little or no alloy elements beyond the Mn, P, S, Si, and Al needed to produce a good quality structural material. The low-alloy steels are the next group that can be characterized by small additions of Cr, Mo, and Ni, usually in the range of about 0.10–4.0% of each element, but generally less than 5% of the total alloying elements. Higher additions of these elements form a group of steels referred to as alloy steels. Generally, the alloying content is equal to or less than 10% (e.g., 9 Cr–1 Mo steel). The distinction between low-alloy and alloy steels is not well defined nor even well observed in practice. Often, all of these steels are lumped together under the term ‘‘low-alloy’’ steel or ‘‘alloy steel.’’ As will be seen in this chapter, 1 2 Craig this distinction is largely unnecessary from a corrosion standpoint because alloying of less than 10% for many environments is not sufficient to impart significant corrosion resistance to steels. In a similar vein, cast irons are used as structural or pressure-containing alloys that have little natural corrosion resistance. Additions of Cr, Ni, and Si are most often the primary means for improving corrosion resistance. Unquestionably the most important alloying element in steels and irons from a corrosion standpoint is Cr. Steels containing in excess of 11% Cr, will display stainless (rust-resistant) qualities when exposed to the atmosphere; thus, this group or, more properly, family of alloys is termed ‘‘stainless steels.’’ In this chapter, the ferritic and martensitic stainless steels will also be discussed. Although the number of alloys that are covered by the categories just presented are myriad, the general performance is relatively easy to address. The corrosion performance of these ferrous alloys is of major importance not only because they represent the largest tonnage of metals used by the world but because they represent the benchmark from which corrosion performance of other alloys is compared. II. CORROSION BEHAVIOR A. General Corrosion Carbon and low-alloy steels generally display active corrosion in the majority of environments to which they are exposed. This means they will corrode unabated at some corrosion rate determined by factors such as solution composition, pH, fluid velocity, presence of oxidizers, temperature, and so forth. In many of the environments to which ferrous alloys are exposed, there is little effect or benefit of minor alloying element additions. Figure 1 illustrates the typical polarization behavior for steels in many environments. The anodic curve shows active corrosion with no tendency toward passivation. The environmental factors mentioned earlier will determine the anodic and cathodic behaviors and ultimately the anodic current density (i.e., the corrosion rate). Figure 2 illustrates the effect of increasing the conductivity of the solution (produced by increasing chloride content) on the corrosion rate of carbon steel (1). Solution conductivity plays a major role in the tendency for corrosion of alloys in a specific environment. For example, in many hydrocarbon environments, the conductivity and polarizability of the solution are so low that corrosion cannot be established or maintained. In these solutions, carbon steel is quite useful and cost-effective. Likewise, in systems containing corrosive gases (i.e., CO2, H2S, etc.), if no water is present, there is no electrolyte for corrosion, and carbon steels are adequate for the service. Ferrous Alloys 3 Fig. 1 Typical polarization behavior for mild steel under active corrosion. In those environments in which corrosion of steel follows the behavior in Fig. 1, other means of mitigating corrosion must be considered. These other means are coatings, inhibitors, cathodic protection, or anodic protection. These other methods are dealt with in more detail elsewhere (2). In some very specific environments, steels may develop a protective corrosion product layer that essentially passivates the steel surface, reducing corrosion to an acceptable level. Although the environments for which this phenomenon occurs are few compared to those for active corrosion, they are notable. Examples of such environments are steels exposed to concentrated sulfuric acid, hydrofluoric acid, and sodium hydroxide. In concentrated sulfuric acid, a soft protective iron sulfate corrosion product is formed that inhibits further corrosion. However, 4 Craig Fig. 2 Conductivity effects on mild steel in aqueous solutions, argon saturated, as a function of NaCl content. this film is not mechanically strong and is easily eroded. Thus, this film is not suitable for exposure to high-velocity streams, yet it is beneficial from a sulfuric acid storage standpoint because carbon steel containers can be used to handle the acid under essentially static conditions. Other environments produce this same behavior and Fig. 3 illustrates this Ferrous Alloys 5 Fig. 3 Passive behavior of mild steel in environments such as Na2SO4. development of passivity in the anodic curve that reflects a decrease in the anodic corrosion current with the formation of a passive film (1). Great care must be taken in applying this method of passivity, however, because many factors in actual service can eliminate or degrade this protective film, causing significant corrosion to occur. Velocity changes, temperature increases, the presence of impurities (i.e., chlorides), and concentration changes can produce high-corrosion rates instead. A useful example of this change in 6 Craig corrosion rate is carbon steel in 90% sulfuric acid at room temperature. Under static conditions, the corrosion rate is about 0.5 mm/year. However, at a concentration of less than 50% H2SO4, the corrosion rate exceeds 5 mm/year. Thus, the stability of the passive film is an important factor in the choice of any material for a specific environment and that choice may be suitable only over a narrow range of conditions. As will be discussed in later chapters, the stability of the passive layer on nickel-based alloys, titanium alloys, and other materials is much greater than for steels, thus the reason these alloys are more resistant to corrosive environments. It is the great stability of the air-formed oxide on ferritic and martensitic stainless steels that produce their stainless quality when exposed to the atmosphere. Yet, this passive film is not stable in all environments and care must be taken in their application, as the oxide is particularly susceptible to attack by halides. B. Localized Corrosion In addition to the uniform or general corrosion of ferrous alloys, there are numerous forms of localized corrosion that can cause failure of these alloys. Pitting corrosion is a highly localized attack of the metal, creating pits of varying depth, width, and number. Pitting may often lead to complete perforation of the metal with little or no general corrosion of the surface. This can be a considerable problem in steels and is one of the most common causes of failure for stainless steels. At this time it is impossible to predict the remaining life of a pitted structure; thus, pitting remains one of the leading causes of failure for ferrous alloys. Crevice corrosion is similar to pitting corrosion in its localized nature but is associated with crevices. Stainless steels and some nickel-based alloys are particularly susceptible to this form of corrosion; however, steels are less susceptible to this form of attack, except in aerated environments. Intergranular corrosion is the preferential corrosion of grain boundaries in a metal caused by prior thermal treatments and related to specific alloy chemistries, especially in stainless steels and nickel alloys. Corrosion of this type is rare in carbon and alloy steels but can be a problem in ferritic and martensitic stainless steels. Dealloying is the selective removal of one element (usually the least noble) from an alloy by the corrosive environment. Also referred to as selective leaching or dezincification, denickelification, and so forth, designating the element removed. Steels are not generally attacked by this mechanism, nor are ferritic or martensitic stainless steels. However, some cast irons, especially gray iron, are quite susceptible to dealloying. For gray cast iron, the graphite flakes are cathodic to the surrounding ferritic matrix. Thus, the ferrite is selectively corroded away, leaving a mechanically weak graphite structure. Corrosion fatigue is the initiation and extension of cracks by the combined action of an alternating stress and a corrosive environment. The introduction of Ferrous Alloys 7 a corrosion environment often eliminates the fatigue limit of a ferrous alloy, creating a finite life regardless of applied stress level. It is currently impossible to predict the corrosion fatigue life of an alloy because of the difficulty in distinguishing the contributing effects of the corrosion portion and the mechanical portion of corrosion fatigue. Galvanic corrosion is the accelerated corrosion of the least noble metal when coupled to one or more other metals. The more noble metals are protected from corrosion by this action. This form of attack is one of the most common causes of corrosion for all of the ferrous alloys, especially carbon and alloy steels. More detail on this type of corrosion is provided later in this chapter. Many forms of flow-assisted corrosion are often included under the term ‘‘erosion–corrosion’’ such as cavitation, impingement, and corrosion–erosion. All of these types of attack are the result of accelerated corrosion due to flow of solids, liquids, or gases, and the ferrous alloys are very susceptible to this form of attack. Therefore, ferrous alloys are quite limited for applications where a corrosive fluid, even one that is mildly corrosive, is combined with rapid flow. Environmental cracking is the initiation and propagation of cracks by the combined action of a corrosive environment and a tensile stress. Typically, under anodic conditions, this form of attack is most often referred to as stress-corrosion cracking (SCC). Generally, susceptibility to cracking increases with increasing temperature, but not every alloy cracks in every environment. This form of corrosion causes significant damage to steels and stainless steels. Another form of cracking is strictly related to hydrogen absorption into ferrous alloys and the resultant cracking. In aqueous environments and in contrast to SCC, this occurs under cathodic conditions. There are numerous forms of damage associated with hydrogen, which are contained under the collective term ‘‘hydrogen damage (HD).’’ For hydrogen embrittlement and hydrogen-stress cracking, tensile stress and hydrogen atoms are necessary to cause failure. However, contrary to SCC, susceptibility is greatest near room temperature. Other terms and forms are hydrogen-induced cracking (HIC), blistering, sulfide-stress cracking (SSC), hydrogen stress-corrosion cracking, hydriding, and hydrogen attack. There are many other terms too numerous to mention. As with SCC, this is a major problem in steels and martensitic stainless steels. Although all of these corrosion mechanisms are of some concern for ferrous alloys, the three most problematic and often observed forms are pitting, galvanic corrosion, and environmental cracking (this term is frequently used to encompass all forms of SCC and HD). Therefore, these three forms will be discussed in greater detail as they relate to ferritic and martensitic steels. 1. Pitting Corrosion Pitting corrosion is one of the most common and most insidious types of corrosion attack on steels. Pitting may rapidly produce perforation of a metal or may take 8 Craig many years to develop. Currently, there are no methods to accurately predict the propagation rates of pits and, therefore, no valid means to estimate the remaining life of a structure or component once pitting has initiated. There has been some success in modeling pitting as a stochastic process, but, as yet, there is not an accepted methodology. Because of this inability to predict pitting and remaining service life, the primary focus in materials selection for a specific environment is to choose a material that is either immune to a particular environment or at least highly resistant to pitting in the first place. This can be a difficult task because often it is not the major component of the service environment that induces pitting but rather the small concentration of some impurity that does. A good example of this is shown in Table 1, where increasing the Cl⫺ content of H 2 SO4 requires a corresponding increase in Cr to the steel to resist pit initiation (3). As discussed earlier, steels are generally resistant to concentrated H2SO4, but the introduction of small amounts of Cl⫺ makes the solution particularly corrosive. This same effect is observed for steels exposed to seawater. Seawater itself is not very aggressive to ferrous alloys; however, it is the introduction of dissolved oxygen that causes seawater to become corrosive, producing severe pitting attack. Figure 4 illustrates the dramatic effect of oxygen, in only the parts-per-billion (ppb) range, on the corrosion of steel (4). The mechanism of pitting is well understood in a general sense. Pit initiation begins with the very localized breakdown of the passive film, leading to the formation of a small pit bottom that acts as the anode and the remainder of the passive surface as the cathode. Thus, there is a large driving force to continue development and propagation of the pit. However, the nature of pitting is a selfsustained autocatalytic process that continues pit propagation. During the propagation process, the solution in the pit bottom becomes and remains very acidic, further enhancing propagation. Moreover, the potential difference between the steel surface and the pit bottom acts as a driving force for propagation. During Table 1 Minimum Concentration of Cl⫺ Necessary for Pit Initiation in 1N H 2SO4 Solution Alloy Fe Fe–5.6 Cr Fe–11.6 Cr 18.6 Cr–9.9 Ni–Fe 20.0 Cr–Fe 24.5 Cr–Fe 29.4 Cr–Fe Cl⫺ (normality) 0.0003 0.017 0.069 0.1 0.1 1.0 1.0 Ferrous Alloys 9 Fig. 4 Effect of oxygen concentration on corrosion of mild steel in Pacific Ocean water. this period, the original steel surface, which has not begun pitting, is effectively protected from further corrosion by the resulting cathodic polarization. The difficulty in predicting the remaining life of a structure during pitting corrosion is due to continual pit initiation, propagation, repassivation, and repropagation. Not all pits in the same structure propagate at the same rate and propagation is not linear but rather an exponential function. It is generally recognized that pitting will initiate at microstructural discontinuities on the steel surface. These discontinuities can be grain boundaries, second-phase particles, and so forth, but they are most often sulfide inclusions. This latter feature is most commonly the origin for pits in stainless steels. Therefore, it is quite predictable that resulfurized steels, especially the resulfurized stainless steels, will suffer pitting corrosion in an environment long before and under less severe conditions than the lower sulfur version of the same steel. For example, AISI 416 stainless steel, which contains 0.15% S minimum, compared 10 Craig to its counterpart, AISI 410, which contains 0.030% S maximum, is highly susceptible to pitting corrosion and may pit in environments where AISI 410 does not. It is impossible to list all the environments in which steels pit because of the numerous factors involved and the great variety of possible combinations of chemicals. Moreover, great care must be taken when selecting an alloy for a certain application in not simply scanning the large number of corrosion data references and databases for alloys with low corrosion rates, as most of these resources do not present pitting data, but rather provide only uniform corrosion rates that can be very misleading. However, that said, it is often the case that steels and, for that matter, many other alloys have a great tendency to pit in environments that contain chlorides or, more generally, halides. Although chlorides are by far the most prevalent species in many environments bromides, iodides and fluorides can also induce pitting. Therefore, the presence of halides in a process stream should be a signal that pitting must be considered in the choice of alloys. Yet, the absence of these species does not necessarily eliminate the possibility of pitting. An example of pitting in the absence of halides is corrosion from CO2 gas dissolved in water. This condition produces carbonic acid that can lead to pitting corrosion of carbon and low-alloy steels. Of course, the situation becomes more complex as a function of temperature and the introduction of chlorides. Figure 5 shows the envelope of applicability of AISI Type 420 stainless steel (also referred to as 13 Cr) to a combined environment of CO2 and chlorides as a function of temperature in the absence of oxygen (5). Within this envelope, no pitting occurs and corrosion is minimal but uniform. However, the introduction of small concentrations (ppb) of oxygen creates a severe pitting attack of the 13 Cr even at ambient temperature, thereby eliminating the use of this alloy. Thus, prior experience or laboratory testing is often necessary to confirm that a particular alloy will not be susceptible to pitting in a specific environment. 2. Environmental Cracking As previously indicated, environmental cracking (EC) is a general term that encompasses all forms of cracking that are induced or accelerated by the service environment. The two principal categories within this form of corrosion that are pertinent to this discussion on ferrous alloys are SCC and HD. An in-depth review of these types of cracking and their mechanisms can be found elsewhere (6). It is simplest and consistent with much of the literature to discuss SCC in terms of an active path corrosion coupled with a tensile stress (often referred to as anodic cracking) and HD as all those forms of cracking that depend on hydrogen assistance (often referred to as cathodic cracking in aqueous environments, Ferrous Alloys 11 Fig. 5 The corrosion resistance of 13 Cr (Type 420) stainless steel in CO2 /NaCl environments in the absence of O2. hydrogen-stress cracking, HIC, blistering, and hydrogen embrittlement, to name a few). In many environments where steels are susceptible to SCC, a frequent precursor to crack initiation is pitting corrosion. In these environments when stresses, either applied or residual, are relatively low, pitting ultimately produces failure. However, as the stress level increases, SCC can become the controlling mode of failure. Figure 6 illustrates this sequence of events (7). The important feature of SCC is that cracking initiates and propagates at a subcritical level below the scale of macroscopic flaws that would be considered critical from a linear elasticfracture mechanics (LEFM) standpoint. Therefore, LEFM by itself cannot be used to predict the likelihood of EC. One of the most significant factors affecting EC is the strength level of the steel. High-strength steels are very susceptible in a variety of environments and this susceptibility is a function of the yield strength. Figure 7 shows that many steels fail in a simple marine environment at ambient temperature when the yield strength exceeds about 180 ksi, regardless of alloy composition (8). However, below 150 ksi yield strength, cracking does not occur in this environment. Figure 12 Craig Fig. 6 Proposed sequence of crack initiation, coalescence, and growth for steels undergoing subcritical cracking in aqueous environments. 7 should not be construed to mean that ferrous alloys do not crack at all below 150 ksi, because quite low-strength steels are very susceptible to EC, just in different environments. Moreover, the cracking of these high-strength steels in seawater is thought to involve HD. Thus, it is convenient to further discuss the EC of ferrous alloys in two groups; high-strength steels (⬎150 ksi) and lowstrength steels (ⱕ150 ksi). a. Environmental Cracking of Low-Strength Steels Low-strength steels (ⱕ150 ksi yield strength) are quite susceptible to EC in certain specific environments. The yield strength of the steel in this strength range is not particularly significant to the susceptibility to EC as it is for higher-strength steels. Rather, other factors such as applied stress, steel composition, pH, solution composition, potential, and temperature are much more critical. Increasing applied (or residual) stress and increasing temperature enhances the SCC of low-strength steels as does decreasing pH. Small concentrations of trace or impurity elements in the alloy can have a profound effect on SCC of steels. Some of the more common environments known to cause SCC of lowstrength steels are liquid NH3, CO2 /CO, carbonate/bicarbonate, hydroxide, nitrate Ferrous Alloys 13 Fig. 7 Stress-corrosion behavior of steels exposed to marine atmosphere. solutions, and amine solutions. Generally, as the concentration of the solution increases, the susceptibility to SCC increases. Table 2 shows the effect of increasing nitrate concentration on the threshold stress of a plastically deformed low carbon–manganese steel (9). The threshold stress (that stress below which cracking does not occur) decreases with increasing concentration of nitrate and is partially dependent on the specific cation associated with the nitrate anion. Similar behavior has been observed for OH solutions and sufficient data have been gathered to develop the useful engineering diagram shown in Fig. 8 Table 2 Threshold Stress Values (ksi) for Mild Steel in Boiling Nitrate Solutions of Various Concentrations Solution concentration Nitrate 8N 4N 2.5N 1N NH 4 NO3 Ca(NO3 ) 2 LiNO3 KNO3 NaNO3 2.2 5.6 5.6 6.7 9.0 3.4 7.8 9.0 4.5 9.5 7.8 13.4 21.3 (2N ) 15.7 24.7 13.4 25.8 25.8 26.9 29.2 14 Craig Fig. 8 Temperature and concentration limits for stress-corrosion cracking susceptibility of carbon steels in caustic soda (NaOH). (10). Area C can also be handled successfully with austenitic stainless steels. This diagram illustrates the important effect of residual welding stresses on SCC and the ability to extend the range of applicability of steels in OH simply by reducing the residual stresses. An area of great concern that has recently received increased attention is the SCC of low-strength pipeline steels. The external SCC of pipeline steels has occurred in two distinct environments. Early failures were in soil environments that produced solutions of carbonate/bicarbonate with a pH of about 9.5 on the outside of the pipe, causing intergranular SCC. Figure 9 shows the intergranular SCC of a low-strength pipeline steel that failed in the high-pH environment. More recently, transgranular SCC has been found to be the cause of several pipeline failures. The pH in this latter case frequently falls in the range of 6–8. Many of Ferrous Alloys 15 Fig. 9 Intergranular fracture of a low-strength pipeline steel from SCC. Magnification: 100⫻. the pipeline failures occurred in pipelines that are more than 20 years old that have yield strengths around 52,000 psi. This long incubation time for crack initiation and propagation is typical for low-strength steels and in sharp contrast to the often rapid initiation and fracture of high-strength steels. Yet, it would be misleading to assume that SCC of low-strength steels is always a slow process. Figure 10 shows that the crack growth rate in many low-strength steels is a strong function of the solution composition and is directly related to the bare surface current density (11). This current density, in combination with straining at the crack tip, is the driving force for cracking and is frequently referred to as active path or anodic cracking. It is generally believed, though not entirely agreed, that SCC progresses by the rupture of the oxide film at the crack tip, thereby providing a bare surface for the peak current to advance the crack tip a certain distance 16 Craig Fig. 10 Crack velocities and peak current densities at the same potentials for a variety of systems and alloys. before the oxide re-forms and the crack arrests. These events may occur over many cycles or just a few. However, it is important to recognize the difference in this mode of cracking versus that due to cathodic cracking, where hydrogen is the primary agent to assist cracking. In general, martensitic stainless steels are used at higher strengths than ferritic stainless steels because the former can be strengthened by heat treatment and the latter cannot. Therefore, the martensitic stainless steels are discussed under the high-strength section. Although ferritic stainless steels are generally more resistant to SCC than austenitic stainless steels, especially in chloride solutions, they are not entirely immune. Small additions of Ni and plastic deformation can each increase the tendency for SCC in chloride environments. Hydrogen damage of low-strength steels typically occurs in steels that have yield strengths less than 100 ksi. As with the SCC of low-strength steels, the yield strength is not an important factor in HD. Moreover, residual and/or applied stresses have little effect. Ferrous Alloys 17 The primary cracking modes are stepwise cracking (SWC, also referred to as HIC) and blistering (HB) or blister cracking. Both types of cracking are the result of relatively high hydrogen input fugacities compared to the HD of highstrength steels and are often found together in the same steel. Both SWC and HB are considered to occur by the classical hydrogen-pressure mechanism. According to this mechanism, hydrogen atoms enter the steel and combine at discontinuities (i.e., nonmetallic inclusions) to form molecular hydrogen, which is too large of a molecule to diffuse back out of the steel. The molecules continue to accumulate, increasing the local hydrogen pressure until a crack or blister forms. Figure 11 shows an example of SWC in a low-strength steel exposed to H2S. Hydrogen damage of steels occurs over the entire strength range of typical engineering applications. Figure 12 shows that regardless of the strength level of the steel, some form of hydrogen cracking may occur and the only distinction is in the morphology of cracking (12). b. Environmental Cracking of High-Strength Steels The EC of highstrength steels (⬎150 ksi) is highly dependent on strength, and in many environments, it is difficult to distinguish between the more classical HD and SCC mech- Fig. 11 Stepwise cracking from hydrogen in a low-strength steel exposed to H2S. Magnification: 25⫻. 18 Craig Fig. 12 Critical hydrogen concentration in steel for cracking as a function of yield strength and the morphology of cracking. anisms. On the other hand, EC from caustic solutions are not obviously hydrogen related, but the crack propagation rate is many times faster than for low-strength steels in caustic. Seawater and brackish waters do not typically produce EC failures of lowstrength steels, but they do produce the EC of high-strength steels, as demonstrated earlier in Fig. 7. Again, this is most likely a HD mechanism. From an engineered materials sense, the actual mechanism is not as important as the fact that high-strength steels are so susceptible to EC, and the resulting crack propagation rate so high that catastrophic failure in many otherwise benign environments can easily occur. Because of this high risk of steel failure with increasing strength beyond 150 ksi, it is common practice to select other alloys and materials that have a greater overall corrosion resistance for high-strength applications. These materials and their performance are dealt with in the remainder of this book. Ferrous Alloys 19 As an example of the behavior of high-strength steels in some of these environments, Fig. 13 shows the crack propagation rate of various alloy steels (13). It is apparent that above 150 ksi yield strength, the crack growth rates are so high that strength level becomes meaningless. Moreover, the speed at which the crack propagates is too rapid for detection in actual service, often leading to a catastrophic failure. Martensitic stainless steels are generally resistant to chloride SCC when heat treated to yield strengths less than about 100 ksi. However, above this yield strength, they become increasingly susceptible to EC in seawater and H2S. Both of these environments are known to produce hydrogen so that the failure of highstrength martensitic stainless steels in these cases is probably a HD mechanism. When selecting an alloy or material for a specific application, it is common practice to first ensure that EC will not be a potential problem in service. Once this form of degradation is eliminated, the select material can be further evaluated for resistance to other less catastrophic forms of attack. Fig. 13 Comparison of stress-corrosion crack velocities in maraging and low-alloy steels. 20 Craig C. Galvanic Corrosion Galvanic corrosion is one of the most common yet least recognized corrosion problems for ferrous alloys. When two or more dissimilar metals are intimately connected and placed in a solution, current will flow between them because of the potential difference of the metals. The metal with the least resistance to corrosion (active metal) in the particular environment will become the anode, and the more corrosion-resistant metal (noble metal) will become the cathode. Corrosion of the anode will usually be more severe than if that metal was alone in the same solution, whereas the cathode will achieve a degree of protection from the environment—sometimes to the extent that corrosion is completely stopped on the cathodic metal. These effects can be measured and have been done for couples of metals in seawater at 25°C. Table 3 provides a relative ranking of metals and alloys regarding their resistance to corrosion in seawater (14). The greater the distance between two metals in the table, the greater their potential difference and the higher the probability that the active metal will suffer accelerated corrosion. Note that some alloys and metals are listed twice in the table: once with the word ‘‘active’’ following and once with the word ‘‘passive.’’ Some metals and alloys become essentially immune to corrosion in certain environments beTable 3 Galvanic Series of Some Commercial Metals and Alloys in Seawater Active or anodic Magnesium Magnesium alloys Zinc Galvanized steel Aluminum 1100 Aluminum 2024 Mild Steel Wrought Iron Cast Iron 13% Chromium stainless steel Type 410 (active) 18-8 Stainless steel Type 304 (active) Lead–tin solders Lead Tin Muntz metal Manganese bronze Naval brass Nickel (active) Nickel Alloy 600 (active) Yellow brass Admiralty brass Red brass Copper Silicon bronze 70–80 Cupro nickel G–bronze Silver solder Nickel (passive) Nickel alloy 600 (passive) 13% Chromium stainless steel Type 410 (passive) 18-8 Stainless Steel Type 304 (passive) Silver Graphite Gold Platinum Noble or cathodic Ferrous Alloys 21 cause of the formation of a surface film so thin that it is impossible to see with the naked eye or even with an optical microscope. The stability of these films is paramount to the enhanced corrosion resistance of these alloys. Moreover, corrosion films represent the controlling factor in almost all corrosion (15). The addition of chromium to iron can produce a passive alloy (18-8 stainless steel) of considerable corrosion resistance compared to the original iron. However, in an environment in which the passive film is not functional, the active surface becomes far less noble, as indicated in Table 3. A more active metal in the series will corrode at the expense of a nobler one. Thus, coupling zinc to steel will cause the zinc to corrode and will protect the steel. This is the reason for galvanizing steel; when pinholes in the galvanizing occur, the steel underneath, once exposed to the environment, will be protected by the zinc. This is also the basis for cathodic protection. Sacrificial anodes are made of metals or alloys that are more active than steel, which allows for the consumption of the anode and the protection of the steel structure. However, if steel is coupled to copper, the distance on the chart is large, so in this case, steel will be the anode and have a greater tendency to corrode. The galvanic series is useful for approximating the behavior of coupled alloys; however, care must be used in its application. Several parameters affect galvanic corrosion and, as such, may affect the actual behavior of a couple in service. Two important factors in galvanic corrosion are the temperature and the relative area of the metals. Increasing temperature in some cases may cause a reversal in the anode–cathode relationship. This reversal has been responsible for failures of galvanized systems or systems protected with zinc sacrificial anodes. These effects point to the need to measure the potential of a couple, especially in cathodic protection, in the actual environment prior to its application. It must always be borne in mind that the ranking of alloys in Table 3 is strictly true only for seawater and that extending it to other environments may result in some changes in the position of metals and alloys in the series. The other factor, the ratio of the area of the anode to the area of cathode, is of considerable importance. If the anode area is smaller than the cathode area, the corrosion rate may be increased many orders of magnitude as a function of this ratio. However, if the anode area is greater than the cathode area, corrosion of the anode will be less than for a 1: 1 anode/cathode ratio. III. EFFECTS OF ALLOYING ELEMENTS Small additions of alloying elements to ferrous alloys generally do not significantly improve their corrosion resistance. As stated earlier, at least 11% Cr is needed to ensure that a steel becomes stainless and thus possesses a certain degree of corrosion resistance. One important exception to this behavior is the class of 22 Craig Fig. 14 Corrosion rate of steels in wet CO2 as a function of the chromium content of the alloy. steels referred to as weathering steels. Small additions (⬍0.5%) of elements such as Cr, Ni, Cu, and P greatly enhance these steels’ resistance to rusting in the atmosphere. This is accomplished through the production of a tight tenacious rust that forms on the steel after exposure to its environment. Hence, further corrosion is stifled. However, for most other service environments, the small variations in alloying or tramp elements are not sufficient to increase the corrosion resistance of a steel. Figure 14 shows the benefit of increasing the Cr content on steels exposed to water containing a high concentration of dissolved CO2 (16). This behavior is typical for steels exposed to many acids as well as other environments. Thus, Cr is an important alloying element for a steel’s resistance to acid Ferrous Alloys 23 attack. A similar benefit is observed for cast irons exposed to nitric and hydrochloric acid by alloying with either Si or Ni in excess of about 5%. IV. EFFECTS OF HEAT TREATMENT Carbon and low-alloy steels and the martensitic stainless steels can be and are heat treated to enhance certain mechanical properties and, as a result, develop different microstructures. Various combinations of ferrite, pearlite, bainite, and martensite may be present in a particular steel, depending on its thermal history. By and large, heat treatment and the subsequent phases formed are not a significant factor in the corrosion of steels. However, there are important exceptions. As previously discussed, environmental cracking is highly dependent on steel strength and, to some degree, its microstructure. Likewise, in certain corrosive environments, a distinction can be observed in corrosion rate as a function of heat treatment. One important example is shown in Fig. 15, for Type 420 stainless steel in wet CO 2 (17). In actual practice, heat treatment/microstructure is not a large concern, except for environmental cracking. V. EFFECTS OF SOLUTION VELOCITY In general, the corrosion rate of ferrous alloys increases with increasing velocity. This can be readily explained by the increased mass transport of ferrous ions across the fluid boundary layer that is established under flowing conditions com- Fig. 15 Effect of the tempering on the yield strength and on the corrosion resistance of quenched-and-tempered 13% Cr stainless steel. 24 Craig bined with the enhanced transport of corrodents to the metal surface across this boundary layer. Of course, the picture is actually more complex when the Helmholtz double layer and the presence of a corrosion product film are included. However, regardless of these issues, at some critical velocity these layers are essentially stripped away and bare metal is continually exposed to the fluid stream. At this point, one of two entirely divergent phenomenon may occur. The erosion–corrosion rate becomes extremely high due to the loss of the ratedetermining mass transport across all of these layers or the erosion–corrosion rate becomes much lower as a result of the inability of the corrodent to reach the bare metal surface and have sufficient time to react. Both of these phenomena are observed for ferrous alloys. Figure 16 illustrates the former case for carbon steel in distilled water (18). At different pHs, the corrosion product formed can be resistant to erosion–corrosion (pH ⬃ 5 and 9), thus minimizing the erosion– corrosion rate, or the corrosion film is unstable (pH ⬍ 4, pH 6–8), leading to high erosion–corrosion rates. The introduction of solid particles such as sand significantly lowers the erosion–corrosion threshold, making the selection of alloys resistant to erosion–corrosion much more difficult. In the absence of solid particles, it has been found that erosion–corrosion resistance is strongly a function of the nature of the oxide. Therefore, more corrosion-resistant alloys such as stainless steels, nickel-based alloys, and titanium alloys have greater erosion– corrosion resistance than steels, even if the alloy is much softer, because of their tighter more resilient oxides. Fig. 16 Effect of pH of distilled water on erosion corrosion of steel at 50°C and 12 m/s flow rate. Ferrous Alloys 25 VI. GENERAL APPLICABILITY OF FERROUS ALLOYS Every service environment is different and care should be taken in trying to generalize the performance of ferrous alloys, especially carbon and low alloys, in common environments. However, several important characteristics are worth emphasizing. In almost every circumstance, the presence of dissolved oxygen in solution causes an increased corrosion rate for steels and cast irons. However, for ferritic and martensitic stainless steels, the presence of oxygen is not as critical and often oxygen in only the ppb range is necessary to maintain the oxide film. Moreover, under total anaerobic conditions at sufficiently high temperatures, oxygen is available from dissociation of the water molecule. The problem for steels and cast iron is due to the fact that oxygen is a very effective cathodic depolarizer; that is, it stimulates the cathodic reaction, and because the anodic and cathodic reactions are interdependent, it produces a net increase in the corrosion rate. Likewise, steels are not resistant to corrosion in acidic pH environments or even in many neutral pH environments, especially in the presence of dissolved oxygen. It is for this reason that steels are most often painted (coated) or cathodically protected to ensure a satisfactory service life. Exposure to the atmosphere is often sufficient to cause corrosion of ferrous alloys to the extent they become either unserviceable or, more typically, aesthetically unpleasing. Moisture, temperature, periods of wet and dry, the presence of chlorides (coastal locations), and industrial pollutants (oxides of sulfur and nitrogen) all contribute to atmospheric corrosion of steels and cast iron. Atmospheric corrosion requires a critical moisture content in the atmosphere and its rate generally increases when the humidity exceeds this critical value (i.e., approximately 60%). Figure 17 illustrates this phenomenon for corrosion of many alloys, including steels, as a function of relative humidity (RH). A form of atmospheric corrosion, wet corrosion, is considered to occur when actual water layers or pools form on the surface of the metal, often from dew, rain, or sea spray. This can be a very complex state because a thin layer of water can act to dissolve a high concentration of gases from the atmosphere, causing a concentrated solution at the metal surface, which produces a correspondingly high, short-term corrosion rate that produces a locally high metal ion concentration, resulting in an oxide that stifles further corrosion or, if the corrosion product is soluble, continued localized attack. Temperature can have many secondary effects on atmospheric corrosion besides the primary effect of increasing reaction rate. Temperature influences the relative humidity, dew point, and time of wetness; all important factors in themselves on atmospheric corrosion. Contaminants, essentially airborne in nature, can profoundly affect atmospheric corrosion. Agents such as gases, like SO2, that can selectively absorb on 26 Fig. 17 Craig Corrosion of ferrous alloys as a function of relative humidity. metal surfaces which act as a catalyst to form SO3 and thus H 2 SO4 in moist environments or particulates such as dust that can cause local cells to form by aiding in the absorption of water and chlorides can accelerate corrosion of alloys that normally would be resistant to atmospheric corrosion in a relatively clean environment. Climatic conditions have a variable effect on corrosion rate. For example, in some regions, winter exposure may be more severe than summer if fuels are used during cold spells that increase combustion products in the air such as NO x and SO2. Conversely, if these fuels are not used in the region, then summer may be worse due to the higher metal temperatures. Likewise, periodic rainfall may be beneficial causing a rinsing action on the surface compared to a climate where the surface is continually wet. Thus, time of wetness and wet/dry cycles can also be quite important, especially if frequent periods of wet and dry can limit the formation and development of a protective oxide layer. Moreover, the existence of insoluble corrosion products can act to entrain water during short dry cycles, keeping the metal surface sufficiently wet to continue corrosion. Two other external environments that can cause significant deterioration of ferrous alloys are soils and concrete. The majority of buried structures in the Ferrous Alloys 27 world are made of cast irons and steels. Soils have a great variability in their tendency to cause corrosion of ferrous alloys. Some of the more important factors that contribute to the aggressiveness of soils are resistivity, pH, moisture, oxygen, bacterial activity, and temperature of the ferrous alloy (i.e., hot pipelines) in contact with the soil. Decreasing resistivity and pH, increasing moisture content, oxygen content, and bacterial activity all enhance corrosion of ferrous alloys in soils. The elevated temperature of the steel surface not only can increase the corrosion rate but also can lead to other forms of more serious attack such as EC. As shown in Fig. 9, the external SCC of pipelines has become a particular problem for those pipelines that operate above ambient temperature and the trend in the future is to operate pipelines at even higher temperatures. Often, coatings and cathodic protection can be used to limit these problems; however, they can also exacerbate them if not properly maintained. One of the major reasons older pipelines have become susceptible to SCC is that the coatings have degraded over time and the cathodic protection systems cannot effectively limit corrosion. Corrosion of steel reinforcing bar in concrete has gained attention due to the widespread problem of the crumbling infrastructure (bridges, highways, buildings, etc.) in many countries. Typically, the steel rebar corrodes as a result of the pH of the cement paste in contact with the steel and the diffusion of chlorides and oxygen into and through the concrete. As corrosion products form on the steel, they represent a larger volume than the original iron in place, thereby spalling and cracking the concrete. It has been found that temperature and relative humidity are important factors in rebar corrosion as well as chlorides and oxygen. Thus, tropical climates that are hotter and more humid than temperate climates would be expected to have a greater problem with rebar corrosion than cooler climates. However, even northern climates have had problems for other reasons; for example, deicing salts used on bridges and roadways to eliminate snow and ice can cause severe rebar corrosion. Several of the methods currently used to fight rebar corrosion are organic coated steel, galvanized steel, stainless-steel rebar, and cathodic protection. As can be appreciated from the foregoing comments, coatings on ferrous alloys are an important means of extending the applicability and service life of these materials. As this is a book on engineered materials, it is beneficial to at least mention the general types of coatings used on ferrous alloys. Coatings can be grouped into four general categories: organic, inorganic, conversion, and metallic. Table 4 lists some of the typical coatings under each of these categories. Under the category for metallic coatings, the specific metal is not listed because many metals can be applied; rather, the process is provided because it will determine which metal can be applied. The selection and use of a particular coating is a function of many factors and care must be taken in selecting the right coating for a ferrous alloy. It must always be borne in mind that a coating is part of a system that 28 Table 4 Craig General Categories of Coatings Organic Coal tars Epoxy Phenolics Alkyds Vinyls Urethanes Acrylics Metallic Galvanizing Plating Ion implantation Cladding Flame spray Chemical vapor deposition Physical vapor deposition Conversion Anodizing Phosphating Chromate Molybdate Inorganic Silicates Ceramics Glass Fig. 18 Coating degradation and corrosion of HY80 in artificial seawater. Ferrous Alloys 29 includes the steel substrate; therefore, if the coating is damaged, corrosion often occurs by different mechanisms than if only the steel were involved. For example, galvanizing (Zn) acts as a sacrificial anode to the underlying steel if the coating is damaged. Thus, the steel substrate is protected against corrosion. However, a more noble coating such as Ni on steel can act as a large cathode, thereby accelerating corrosion at a damaged location (referred to as a holiday). Even organic coatings can display accelerated corrosion at holidays in the coating depending on the particular environment to which they are exposed (Fig. 18) (19,20). Generally, coatings show a slow decrease over time in the coating (film) resistance, indicating the gradual permeation of water and other ionic species through the coating. Thus, coating life in any environment is finite and organic coatings are not truly barriers, as is so often mistakenly suggested. However, the application of the correct coating can often double or triple the useful life of a ferrous alloy in certain environments and is, therefore, an important factor in the selection of ferrous alloys. In conclusion, even though ferrous alloys (ferritic and martensitic) are the most widely used engineered materials, they are also the least corrosion resistant—readily degrading in most environments. Because of this behavior, they most often require additional means of controlling corrosion (i.e., coatings, cathodic protection, inhibitors) to provide a satisfactory service life. REFERENCES 1. RL Martin. Application of Electrochemical Polarization to Corrosion Problems. St. Louis, MO: Petrolite Corp., 1977. 2. ASM Metals Handbook, Vol. 13, Corrosion. ASM International, Materials Park, OH, 1989. 3. ND Stolica. Pitting corrosion on Fe–Cr and Fe–Cr–Ni Alloys. Corrosion Sci 9: 455–460, 1969. 4. D Wheeler. Treating and monitoring 450,000 b/d injection water. Petrol Eng Int 1975; November: Vol. 38, 68–72. 5. BD Craig. Selection guidelines for corrosion resistant alloys in the oil and gas industry. Technical Publication No. 10073, The Nickel Development Institute, Toronto, 1995. 6. BD Craig, RH Jones. Environmentally induced cracking. In: ASM Metals Handbook, Vol. 13, Corrosion, ASM International, Materials Park, OH, 1989, pp. 145– 171. 7. FP Ford. Quantitative prediction of environmentally assisted cracking. Corrosion 51:375–395, 1996. 8. EH Phelps. Stress corrosion behavior of high yield strength steels. Proc. Seventh World Petroleum Congress. Amsterdam: Elsevier, 1967. 9. RN Parkins, R User. The effect of nitrate solutions in producing stress corrosion 30 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. Craig cracking in mild steel. First International Congress on Metallic Corrosion, London, 1961, p. 289. Corrosion Data Survey—Metal Selection, 6th ed. Houston, TX: NACE International, 1985, p. 176. RN Parkins. Predictive approaches to stress corrosion cracking failure. Corrosion Sci 20:147–166, 1980. E Sato, M Hashimoto, T Murata. Corrosion of Steels in a Wet H2S and CO2 Environment. Second Asian Pacific Corrosion Control Conf., Kuala Lumpur, 1981. CS Carter. Fracture toughness and stress corrosion characteristics of a high strength maraging steel. Met Trans 2:1621–1625, 1971. Corrosion Basics. LS Van Delinder, ed. Houston, TX: NACE International, 1984, p. 35. BD Craig. Fundamental Aspects of Corrosion Films in Corrosion Science. New York: Plenum Press, 1991. NG Galindez Ruiz. The effect of crude oil on corrosion of alloys in H2S/CO2 environments. Master’s thesis, Colorado School of Mines, 1993. JL Crolet. Acid corrosion in wells (CO2, H2S): Metallurgical aspects. J Petrol Technol 35:1552–1558, August 1983. MG Fontana. Corrosion Engineering, 3rd ed., New York: McGraw-Hill, 1986, p. 94. BD Craig, DL Olson. Corrosion at a holiday in an organic coated-metal substrate system. Corrosion 32:316–321, 1976. JR Scully. Evaluation or organic coating deterioration and substrate corrosion in seawater using electrochemical impedence measurements. Corrosion/86, NACE, Houston, TX, 1986. 2 Austenitic Stainless Steels Russell H. Jones, Stephen M. Bruemmer, and Mike J. Danielson Pacific Northwest National Laboratory, Richland, Washington Bruce Craig MetCorr, Denver, Colorado I. INTRODUCTION Austenitic stainless steels derive their ‘‘stainless’’ properties from the presence of a very effective passive film. This film forms spontaneously within aqueous environments and the stability of this film in various environments determines the corrosion resistance of ‘‘stainless’’ steels. Alloy composition is also very significant in determining the ‘‘stainless’’ character of this family of alloys. Chromium concentration is the one element that directly affects the passive film stability, and Ni, Mn, Mo, C, and N also play a role. Passive behavior begins at about 10% Cr, with increasing film stability occurring with increasing Cr concentration. Repassivation can occur in aqueous environments and it is the rate at which a break in the passive film re-forms that also contributes to the ‘‘stainless’’ character of these materials. Under selected conditions, the passive film is not stable and this can lead to general or localized corrosion phenomena. Austenitic stainless steels are often prepassivated in an acid bath, but the removal of surface contaminants that would hinder passivation in aqueous solutions is the primary purpose of this prepassivation treatment. Given a clean surface in an alloy with 31 32 Jones et al. sufficient Cr, austenitic stainless steels will passivate upon immersion in an aqueous environment. Also, there will always be an air-formed oxide even without this prepassivating treatment. Austenitic stainless steels undergo all the common forms of corrosion, including (1) general, (2) galvanic, (3) pitting, (4) crevice, (5) intergranular, and (6) stress corrosion. General corrosion occurs when ‘‘stainless’’ steel is immersed in an environment in which the passive film is not stable, such as hot sulfuric acid, boiling MgCl2, or another very aggressive environment. Pitting corrosion occurs, as in other metals, because of a local discontinuity in the passive film such as an inclusion or local chemistry change. Halogen ions are the most prevalent cause of pitting in stainless steels, with chloride being the most common halogen to initiate pitting. Pit growth depends on a variety of factors, with the localized corrosion environment in the pit being the most significant factor. Surface condition can also contribute to pitting, with the presence of deposits being a significant factor. Prepassivating treatments are used to clean the surface of deposits. Intergranular corrosion and intergranular stress-corrosion cracking are related phenomena which occur when the grain-boundary microchemistry is altered by thermal treatment such as welding or heat treatment. The process that alters the grain-boundary corrosion resistance is called sensitization and occurs when chromium carbides precipitate at the grain boundaries. II. CORROSION BEHAVIOR A. Alloy Classification Stainless steels are classified by the phases that are present. Table 1 shows the classifications and compositions of the more commonly available alloys, but this table is not exhaustive. The five classifications are austenitic, ferritic, duplex (containing both austenite and ferrite), martensitic, and precipitation hardening (PH). Alloying affects the predominant phase that is present, and the phase has a profound effect on the mechanical and corrosion characteristics. Often, the mechanical properties are the driving force for choosing the alloy for the application, and the corrosion properties must then be optimized within that alloy classification. A brief description of each classification is given in the following subsections. 1. Austenitic Stainless Steels Austenitic stainless steels have a face-centered-cubic (fcc) crystal structure. The most commonly used stainless steels, the 300 series, belong to this group. The austenite is a high-temperature phase that is stabilized by the addition of nickel, manganese, or nitrogen, but they also contain significant amounts of chromium, which gives them good overall corrosion properties. These alloys are fairly low Austenitic Stainless Steels 33 in strength, but they exhibit a high fracture toughness and are used over a wide range of temperatures. As a class, they have good resistance to hydrogen embrittlement. In recent years, a new class has emerged called superaustenitics. They are very high in Mo (4–7%) and nickel, which results in the highest resistance (within the austenitics) to any localized attack processes in chloride-containing media. 2. Ferritic Stainless Steels Ferritic stainless steels have the body-centered-cubic (bcc) crystal structure, using chromium as the major alloying element. They can have higher strengths than the austenitic steels at ambient temperatures, but they suffer from a lower fracture toughness, particularly at lower temperatures. In an attempt to improve the fracture toughness and corrosion behavior, a new class of superferritics was developed. These materials are low in carbon and higher in Cr and Mo than the older ferritics. Successful use of ferritics requires careful attention to detail in controlling the heat treatment. 3. Duplex Stainless Steels Duplex stainless steels have lower amounts of nickel than the austenitic grades, with the result that some of the austenite transforms to ferrite. Generally, the alloying element and heat treatments are controlled to form equal amounts of ferrite and austenite. The principal advantages over the fully austenitic grades are higher strength, improved resistance to stress-corrosion cracking (SCC), and a high immunity to sensitization. In order to improve the localized corrosion behavior, a superduplex series of alloys has emerged. These are alloyed to contain larger amounts of Cr, Mo, and N. 4. Martensitic Stainless Steels Martensitic stainless steels contain significant amounts of chromium (10–18%) and carbon but are low in Ni. Although austenitic at high temperatures, they can be transformed into the martensite structure by rapid cooling. These alloys are very strong, but they suffer a loss of fracture toughness and have generally inferior localized corrosion resistance. Hydrogen embrittlement has been identified as a cause of fracture. Recently, supermartensitics have been formulated that are higher in Mo to improve localized attack behavior. 5. Precipitation-Hardening Stainless Steel Precipitation-hardened stainless steels superficially resemble the 300 series austenitic steels in nickel and chromium composition, but, in addition, they contain small amounts of copper, aluminum, or titanium that can be precipitated with Table 1 Stainless Steel Compositions (wt%) Austenitics 201 202 301 302 304 304L 304LN 316 316LN 316L 321 347 348 384 Superaustenitics 254 SMO AL-6X AL-6XN 20Cb-3 442 446 C Mn Si P Ni Cr S S20100 S20200 S30100 S30200 S30400 S30403 S30453 S31600 S31653 S31603 S32100 S34700 S34800 0.15 0.15 0.15 0.15 0.08 0.03 0.03 0.08 0.03 0.03 0.08 0.08 0.08 5.5–7.5 7.5–10.0 2 2 2 2 2 2 2 2 2 2 2 1 1 1 1 1 1 1 1 1 1 1 1 1 16.0–18.0 17.0–19.0 16.0–18.0 17.0–19.0 18.0–20.0 18.0–20.0 18.0–20.0 16.0–18.0 16.0–18.0 16.0–18.0 17.0–19.0 17.0–19.0 17.0–19.0 3.5–5.5 4.0–6.0 6.0–8.0 8.0–10.0 8.0–10.5 8.0–12.0 8.0–12.0 10.0–14.0 10.0–14.0 10.0–14.0 9.0–12.0 9.0–13.0 9.0–13.0 0.06 0.06 0.045 0.045 0.045 0.045 0.045 0.045 0.045 0.045 0.045 0.045 0.045 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 0.03 S38400 0.08 2 1 15.0–17.0 17.0–19.0 0.045 0.03 S31254 0.02 1 0.8 19.50–20.50 17.50–18.50 0.03 N08366 0.035 N08367 0.03 N08020 0.07 2 2 2 1 1 1 20.0–22.0 20.0–22.0 19.0–21.0 23.5–25.5 0.03 23.50–25.50 0.04 32.0–38.0 0.045 0.03 0.03 0.035 N08904 0.02 2 1 19.0–23.0 23.0–28.0 0.035 S40500 S40900 S42900 S43000 S43600 S43035 0.08 0.08 0.12 0.12 0.12 0.07 1 1 1 1 1 1 1 1 1 1 1 1 11.5–14.5 10.5–11.75 14.0–16.0 16.0–18.0 16.0–18.0 17.0–19.0 — 0.5 — — — 0.5 0.04 0.045 0.04 0.04 0.04 0.04 S44200 S44600 0.2 0.2 1 1.5 1 1 18.0–23.0 23.0–27.0 — — 0.04 0.04 0.045 0.01 Other 0.25 N 0.25 N — — — — 0.10–0.16 N 2.0–3.0 Mo 2.0–3.0 Mo; 2.0–3.0 Mo 5 ⫻ % C min Ti 10 ⫻ % C min Nb 0.2 Co; 10 ⫻ % C min Nb; 0.10 Ta — 6.00–6.50 Mo; 0.50–1.00 Cu; 0.180–0.220 N 6.0–7.0 Mo 6.00–7.00 Mo; 0.18–0.25 N 2.0–3.0 Mo; 3.0–4.0 Cu; 8 ⫻ % C min to 1.00 max Nb 4.0–5.0 Mo; 1.0–2.0 Cu 0.03 0.10–0.30 Al 0.045 6 ⫻ % C min–0.75 max Ti 0.03 — 0.03 — 0.03 0.03 0.15 Al; 12 ⫻ % C min– 1.10 Ti 0.03 — 0.03 0.25 N Jones et al. 904L Ferritics 405 409 429 430 436 439 UNS 34 Type S44660 0.025 1 1 25.0–27.0 1.5–3.5 0.04 0.03 AL 29-4C S44735 0.03 1 1 28.0–30.0 1 0.04 0.03 AL 29-4-2 S44800 0.01 0.3 0.2 28.0–30.0 2.0–2.5 0.025 0.02 Duplex 329 Uranus 50 S32900 S32404 0.2 0.04 1 2 0.75 1 23.0–28.0 20.5–22.5 2.50–5.00 5.5–8.5 0.04 0.03 0.03 0.01 S32550 0.04 1.5 1 24.0–27.0 4.50–6.50 0.04 0.03 S40300 S41000 S41400 S41600 0.15 0.15 0.15 0.15 1 1 1 1.25 0.5 1 1 1 11.5–13.0 11.5–13.5 11.5–13.5 12.0–14.0 — — 1.25–2.50 — 0.04 0.04 0.04 0.06 S42000 S42200 0.15 min 0.20–0.25 1 1 1 0.75 12.0–14.0 11.5–13.5 — 0.5–1.0 0.04 0.04 440A S44002 440B S44003 440C S44004 Lapelloy S42300 Precipitation Hardening 13–8 Mo S13800 0.60–0.75 0.75–0.95 0.95–1.20 0.27–0.32 1 1 1 0.95–1.35 1 1 1 0.5 16.0–18.0 16.0–18.0 16.0–18.0 11.0–12.0 — — — 0.5 0.05 0.2 0.1 12.25–13.25 7.5–8.5 0.07 0.07 0.09 0.07–0.11 1 1 1 0.5–1.25 1 1 1 0.5 14.0–15.5 15.5–17.5 16.0–18.0 16.0–17.0 Ferralium 255 Martensitic 403 410 414 416 420 422 S15500 S17400 S17700 S35000 3.5–5.5 3.0–5.0 6.5–7.75 4.0–5.0 1.00–2.00 Mo 2.0–3.0 Mo; 1.0–2.0 Cu; 0.20 N 2.00–4.00 Mo; 1.50–2.50 Cu; 0.10–0.25 N 0.03 — 0.03 — 0.03 — 0.15 0.6 Mo(b) min 0.03 — 0.03 0.75–1.25 Mo; 0.75–1.25 W; 0.15–0.3 V 0.03 0.75 Mo 0.03 0.75 Mo 0.03 0.75 Mo 0.025 2.5–3.0 Mo; 0.2–0.3 V 0.01 0.008 0.04 0.04 0.04 0.04 0.03 0.03 0.04 0.03 2.0–2.5 Mo; 0.90–1.35 Al; 0.01 N 2.5–4.5 Cu; 0.15–0.45 Nb 3.0–5.0 Cu; 0.15–0.45 Nb 0.75–1.5 Al 2.5–3.25 Mo; 0.07–0.13 N 35 15-5 PH 17-4 PH 17-7 PH AM-350 (Type 633) 0.04 0.04 0.04 0.025 2.5–3.5 Mo; 0.2 ⫹ 4 (% C ⫹ % N) min to 0.8 max (Ti ⫹ Nb); 0.035 N 3.60–4.20 Mo; 0.20–1.00 Ti ⫹ Nb and 6 (% C ⫹ % N) min Ti ⫹ Nb; 0.045 N 3.5–4.2 Mo; 0.15 Cu; 0.02 N; 0.025 max (% C ⫹ % N) Austenitic Stainless Steels Superferritics Sea-Cure (SC-1) 36 Jones et al. the appropriate heat treatment. These alloys can develop extremely high strength and can also share some of the good corrosion properties of the austenitics, but hydrogen embrittlement has been identified as a cause of fracture. This grade of alloy can be martensitic or austenitic. These five classes of alloys can also be obtained as cast alloys. Microstructure is a key variable in controlling the strength, localized corrosion behavior, and SCC behavior of all these classes. It needs to be pointed out that the existence of cast microstructures increase the complexity of choosing the ideal material and making certain the microstructure is under control. B. Composition Effects on Corrosion Stainless steels are vulnerable to all forms of corrosive attack. Their successful use requires (1) knowledge of the alloying constituents that impact resistance, (2) detailed information on the chemical composition and temperature of the test environment, and (3) knowledge of the failure processes to which the particular alloy or alloy class is subject. Each of the major alloying elements will be briefly described to give the reader a general knowledge of how to fit the alloy composition to the environment. 1. Chromium Chromium is the single most important element contributing to the ‘‘stainless’’ behavior of these iron-based alloys, and in general, the higher the chromium level (once it gets above 10%), the better the performance. Additions of chromium greatly improve the behavior over that of iron in neutral and acidic pH ranges but, curiously, has little effect in high-pH environments. The chromium appears in the oxide film and acts to inhibit the transport of corrosion products across it, leading to the formation of a passive film. In particular, the chromium acts to decrease the general corrosion rate and improve crevice corrosion resistance. 2. Molybdenum Molybdenum is perhaps the second most important alloying element from a corrosion standpoint. Small additions have a profound effect by enhancing the passive character of the passive film, particularly in chlorides and reduced sulfur environments. In particular, the pitting and crevice corrosion behavior are improved relative to similar alloys without the molybdenum. Every class of stainless-steel alloy has members that are very high in Mo, and these are called ‘‘super’’ austenitics, martensitics, and so forth. The alloys with high Mo levels have the disadvantage in that certain undesirable phases (sigma, chi, laves) can form unless care is taken in heat treatment and welding. Austenitic Stainless Steels 37 3. Nickel Nickel is the most important austenite stabilizer, but it has a complex effect on the corrosion behavior. It acts to decrease the general corrosion rate in reducing acidic environments. In sufficiently high levels, it acts to improve the SCC behavior in chloride and caustic solutions, but at intermediate levels, it can decrease the SCC resistance. The austenite structure results in a high resistance to hydrogen embrittlement. 4. Manganese At the low levels used in most stainless steels, manganese can be considered a substitute for nickel. It controls the solubility of sulfur by precipitating the sulfur as MnS inclusions. Modern steel-making practice is to keep the sulfur as low as possible because localized corrosion events such as pitting initiate at MnS inclusions. 5. Carbon Carbon is used at low levels as a strengthener in stainless steels, but it also can render the stainless steels vulnerable to localized attack or SCC from sensitization. In sensitization, the carbon reacts with chromium to precipitate a Cr 23C 6 carbide causing a very local decrease of the chromium concentration in the metal, making this depleted region vulnerable to localized attack or SCC. In the austenitics, this has led to the ‘‘L’’ grades, which are low in carbon (less than approximately 0.03 at.%) and fairly immune to this problem. 6. Nitrogen Nitrogen is used at very low levels and acts as a strengthener and austenite stabilizer. Nitrogen greatly improves the pitting and crevice corrosion performance of the austenitics and superaustenitics, particularly in concert with Mo. It helps prevent (slows down) the undesirable chi (and probably sigma, laves) phase from forming during welding or heat-treating operations with the high-Mo alloys. Its behavior is mixed in other classes of stainless steels. C. Corrosion Behavior Choosing the right material is a multistage process. Once the mechanical properties and certain other properties (e.g., wear resistance, weldability, machinability, availability in the needed form, etc.) of the metal are defined for the application, the next level requires a careful definition of the chemical (e.g., pH, chloride level, sulfide level, temperature, velocity, oxygen, hydrogen gas) and exposure environment (single or multiple phases present, occasional dryout, differential 38 Jones et al. aeration, etc.). Once this multistage process is completed, the general class of alloy can be defined. Because cost is always an important parameter for the material choice, the tendency will be to use the lowest alloyed material because nickel and chromium are expensive. As a general rule for iron-based stainless steels, the higher the alloy is in nickel and chromium (and molybdenum), the more corrosion resistant the material will be. The task of the materials scientist is to find the least expensive material that will perform adequately. Often, the exact answer will not be found in the literature; consequently, testing of various representative alloys will be needed. Many test methods are described in Ref. 1. From their inception, stainless steels were developed to have low general corrosion rates. This simplification is still fairly true as long as one avoids certain environments such as reduced sulfur environments or highly acid environments—only highly alloyed stainless steels will perform in these environments. The surprises with stainless steels are usually due to localized attack such as pitting, crevice corrosion, and SCC. Because the chemical environment has infinite variability, the corrosion behavior for each class of stainless steel will be examined using certain defined environments which will act as benchmarks or models. Examples will be potable water (low chloride), seawater (high chloride), and so forth. The book by Sedriks (2) is a particularly good reference. 1. Atmospheric Environments Atmospheric environments are those exposed to the natural elements in the air at ambient temperatures. City and marine environments present the most difficult environments because they are contaminated with sulfur compounds and chlorides. The austenitics, particularly those containing molybdenum, give the best outdoor performance from the standpoint of pitting and crevice corrosion, but it should be noted that they will not be pit-free. Rather, the pit density and depth will usually remain low enough that the cosmetic use of the material will not be impacted. Localized corrosion rates for all classes of stainless steels are highest where the chlorides and sulfides are highest. The ferritics have also been used in this environment but not as successfully as the austenitics. Both the austenitics and ferritics have good immunity to SCC. The higher-strength alloys such as the martensitics and precipitation-hardened alloys are prone to pitting and SCC. This environment can get extremely aggressive if there is a source of heat such that condensation and dryout takes place, resulting in very concentrated solutions. One extreme of this environment might be found under insulation, and, here, SCC and pitting can occur with all the standard stainless classes, even the austenitics and ferritics. The high-Mo (super types) alloys provide the greatest resistance to pitting and SCC when the atmospheric environment is the most aggressive. Austenitic Stainless Steels 39 2. Deionized (Pure) Water (Ambient Temperature) Pure water with low concentrations of dissolved salts (although it may contain dissolved carbon dioxide from the air) is the most ideal environment for stainless steels. They should all perform well and be free of any localized corrosion problem and SCC. 3. Deionized (Pure) Water (High Temperature) The nuclear power industry successfully uses austenitic stainless steels under elevated temperature conditions. Initially, 304 stainless (high C) was used, which when welded gave rise to sensitized grain boundaries which were very prone to SCC. This was a very major problem which has been solved by using low-carbon austenitics. Some ferritics and precipitation-hardened alloys have been successfully used in other components such as turbine blades. The key to successful use is attention to heat treatment and maintaining the purity of the chemical environment. The other classes of stainless steel have not been used extensively in this environment. There is an extensive literature on this subject and the reader is directed to reviews by Hanninen (3) and Cragnolino (4). 4. Fresh (Potable) Water Fresh water can contain a few hundred ppm chloride, and when used in a heat exchanger, it can reach temperatures near 100°C. The 300 series austenitics have been used successfully in this environment, but as the temperature and chloride level increase, the amount of alloying (particularly Mo) must increase to prevent pitting and SCC. Sensitization is an issue and the low carbon grades must be used. For years, there was a widely accepted belief that SCC did not occur below 60°C, but long-term experience has now demonstrated that there is no threshold temperature or chloride level for the 300 series austenitics (5). Consequently, there will be a large variability in the observations of success with the 300 series. The superaustenitics with high levels of molybdenum should perform very well against pitting and SCC degradation. The standard ferritics are low in molybdenum and will be prone to pitting, but they can be resistant to SCC. The behavior of the duplex stainless steels is complex and must be considered carefully. They are more resistant to SCC relative to the 300 series of austenitics but may be similarly prone to pitting. Both the martensitic and precipitation-hardening grades are prone to pitting and SCC. The SCC process for these higher-strength steels is considered to actually be hydrogen embrittlement. 5. Seawater Seawater is the most challenging environment for stainless steels, and the problems are compounded if reduced sulfur species are present due to the action of 40 Jones et al. bacteria or decaying vegetation. In general, the standard alloys within each grade will be prone to pitting and SCC with much variability in the reported results. For maximum safety and reliability, only the high-Mo alloys should be considered for this environment. A review paper by Streicher (6) on 30- and 60-day ambienttemperature crevice tests clearly shows only the superaustenitics, superferritics, high-Mo duplex, and nickel-based alloys are suitable. All localized corrosion and SCC problems worsen at elevated temperatures. Here, the famous boiling magnesium chloride test can give insight into a material’s behavior. The 300 series alloys have completely unsuitable SCC behavior in this test (unless they are cathodically protected). The nickel level has to be above 20% before the austenitics show significant improvement in SCC behavior in the boiling magnesium chloride test. Molybdenum is also beneficial in raising the threshold stress intensity for SCC. Austenitics (high Ni), ferritics, and duplex alloys all can show resistance to SCC, but it is clear that they must be high in Mo. For the most difficult applications, titanium or nickel-based alloys will need to be used. 6. Acidic Environments Chromium is the most important element that imparts resistance to acidic environments; consequently, the highest resistance is associated with the highest Cr levels within each grade of stainless steel. In general, the 300 series austenitic stainless steels can be used in nitric acid over a wide concentration range (0– 65%), even up to the boiling point. However, stainless steels are completely unsuitable for the HCl environment, and mixed acids containing HCl and other halides are also very problematic. The alloy 20Cb-3, containing copper, was especially developed for use in sulfuric acid. Clearly, it is important to define the acid compositions and temperatures and then to utilize the literature for alloy recommendations. General corrosion and localized attack due to microstructural problems are the major causes of failure. Much less is known about the SCC behavior. Organic acids are less corrosive than the mineral acids. 7. Basic Environments In general, all stainless steels are quite resistant to general corrosion in concentrated caustic solutions, even up to boiling temperatures. However, there is a severe SCC problem, and the alloys can act in a very brittle manner. Alloys with the highest nickel content are the most resistant to general attack and SCC. The threshold for SCC is a function of the temperature, caustic concentration, and nickel content. The austenitics 304 and 316 are resistant to SCC up to about 60°C in 60% concentrated caustic. At higher temperatures, nickel-based alloys must be used. Microstructural effects are very important because sensitization makes the materials more susceptible to SCC. Austenitic structures are more resistant than ferritic structures. Austenitic Stainless Steels 41 III. INTERGRANULAR STRESS-CORROSION CRACKING Grain-boundary composition has been inferred to control intergranular (IG) fracture in a wide range of materials systems. Although many authors have attempted to link grain-boundary composition and environmental cracking susceptibility, few have made direct measurements. In most cases, bulk composition and/or heat treatment is varied and it is assumed that interfacial segregation is systematically changed. Indirect measurements are often made (e.g., IG corrosion tests) indicating the grain-boundary composition of an isolated element. Within selected wellunderstood cases, such approaches can give reproducible results. However, quantitative measurements of grain-boundary composition are essential to enable any reasonable assessment of variables controlling cracking susceptibility. With the commonplace use of high-resolution techniques such as analytical transmission electron microscopy (ATEM) and scanning Auger microscopy (SAM), quantitative relationships have been established between interfacial composition and cracking susceptibility for many metallic alloy systems (7). Perhaps the alloy system that has been most closely examined has been austenitic stainless steel due to its widespread use as a corrosion-resistant structural alloy in nuclear power systems. The vast majority of failures have been in high-carbon, 300-series stainless steels thermally sensitized during fabrication. Extensive basic and applied research activities were initiated about 25 years ago to develop a mechanistic understanding of the IGSCC process and, more importantly, to identify remedial actions and corrective measures to cracking problems in boiling-water reactor (BWR) power plants. For the most part, those research activities were highly successful. IGSCC of sensitized stainless steel is probably the best understood and effectively modeled environmental cracking process (8). However, recent observations of IG cracking in cold-worked or in irradiated stainless steels, have been difficult to explain. Austenitic stainless steels provide an example alloy system to demonstrate the influence of grain-boundary composition on IGSCC. Emphasis is placed on identifying equilibrium and nonequilibrium segregants that may promote susceptibility or improve resistance to cracking, respectively. In each case, current understanding of grain-boundary composition development in stainless steels is reviewed and assessed relative to IG fracture in corrosive environments. A. Grain-Boundary Composition and IGSCC The general conditions necessary to promote IGSCC are a susceptible material microstructure–microchemistry, a sufficiently corrosive environment, and the presence of tensile stresses. Many of the important aspects controlling environmental crack advance are illustrated in Fig. 1. In nearly all cases of IG cracking, grain-boundary composition plays a dominant role. Interfacial composition can 42 Fig. 1 Jones et al. Schematic illustrating intergranular stress corrosion processes. be significantly changed from the matrix by equilibrium (nonequilibrium) processes resulting in segregation (depletion) of alloying (impurity) elements and precipitation of second phases. These compositional changes in the grain-boundary region can influence IG crack advance through effects on electrochemical behavior (e.g., dissolution, repassivation, and hydrogen recombination) as well as effects on interfacial mechanical behavior (e.g., deformation and cohesive strength). B. Precipitation and Grain-Boundary Composition Changes The dominant material variable controlling IGSCC susceptibility in austenitic stainless steels results from the precipitation of Cr-rich M23C6 carbides at highenergy interfaces. This promotes the development of a Cr-depleted region adjacent to carbide precipitates. This depletion is controlled by the thermodynamics of carbide formation and differences between the diffusivities of Cr and C. ATEM–EDS has enabled Cr-depletion profiles to be routinely measured demonstrating that interfacial Cr concentrations decrease (from ⬃18% to ⬍10%) as the heat-treatment temperature is decreased due to changes in C and Cr activities. The width of the depleted zone increases with time after IG carbides are nucleated. The extent of grain-boundary Cr depletion has been directly linked to the IG corrosion and SCC susceptibility of austenitic stainless steels (9–12). The threshold concentration to promote IG degradation can be quite different for cor- Austenitic Stainless Steels 43 Fig. 2 Grain-boundary Cr concentration width on IGSCC in BWR water environment. rosion and SCC, as illustrated in Fig. 2. Classical IG corrosion is detected in a standard sensitization test when the grain-boundary Cr concentration drops below ⬃13.5 wt%. On the other hand, IGSCC in high-temperature aerated-water environments can be initiated during slow-strain-rate (SSR) tests when local Cr levels drop below ⬃17 wt%. Additional tests varying the width of the Cr-depletion zone (and keeping boundary Cr concentration approximately constant) reveal that only a very narrow (⬍4 nm) width is necessary to promote cracking (Fig. 3). Fig. 3 Grain-boundary Cr depletion on IGSCC in BWR water environment. 44 Jones et al. IGSCC susceptibility is not sensitive to increases in depletion width beyond that necessary to establish a continuous path for crack advance. On the other hand, standard sensitization tests will show much more aggressive IG corrosion with increasing depletion widths (11). The correlations presented in Fig. 2 point out the critical importance of Cr depletion, and Cr minimums in particular, on IG degradation of austenitic stainless steels. Specific relationships between IGSCC and grain-boundary composition will always depend on many other critical factors, including mechanical loading characteristics and environmental conditions, as well as secondary material variables. For example, the threshold grain-boundary Cr concentration has been shown to depend on the strain rate during SSR tests (10,11) C. Equilibrium Impurity Segregation Impurity elements present at low levels in austenitic stainless steels can reach high levels at grain boundaries due to equilibrium segregation. The most prevalent segregant is P, which can reach grain-boundary P contents ⬎10 at% in commercial stainless steels after intermediate temperature heat treatments (500– 750°C). Thus, materials in the ‘‘sensitized’’ condition will most likely have considerable P segregation along with M23C6 carbides and Cr depletion defining the local microchemistry. Segregation of other impurity elements to stainless-steel grain boundaries has been observed, but this often requires high bulk contents or special thermal treatments. Sulfur segregates rapidly to boundaries if preexisting sulfides are dissolved by a high-temperature (⬎1200°C) exposure. Without such treatment, grain-boundary S segregation is very slight even in doped alloys. Another element that has been shown to strongly segregate to austenitic stainlesssteel grain boundaries is N. Grain-boundary segregation of N in commercial 304 and 316 grades is likely because bulk N levels are typically greater than 0.02 wt% (higher in L grades). Grain-boundary impurity segregation has been shown to promote hydrogen-induced cracking (HIC) in many iron- and nickel-based alloys (13). Phosphorus segregation induces HIC during low-temperature SSR tests at cathodic potentials, as illustrated in Fig. 4, but appears to have no effect on IGSCC in high-temperature water, as indicated by the triangular points in Fig. 2. The influence of grain-boundary Cr content on IG cracking is not affected by P segregation (⬃10 at%). These results are consistent with the crack-growth-rate tests of Andresen and Briant (14), who found that grain-boundary P (and N) enrichment did not promote IG cracking of 304L SS in high-temperature water, whereas S had a small detrimental effect. However, there remains a need for additional crack growth experiments to evaluate segregation effects on SCC in stainless steels strengthened by cold work where cracking has been identified in laboratory tests and in service without Cr depletion (15,16). Austenitic Stainless Steels 45 Fig. 4 Grain-boundary P concentration on HIC in low-temperature acidic environment at cathodic electrochemical potential. IV. NONEQUILIBRIUM THERMAL SEGREGATION Recent ATEM characterizations (17–21) in annealed stainless steels have clearly demonstrated that significant grain-boundary segregation occurs as a result of high-temperature heat treatment and subsequent rapid cooling. Alloying elements that are enriched are Cr and Mo compensated by Ni and Fe depletion, as illustrated in Fig. 5 for 316SS. Normal rapid air cooling from the solution anneal temperature (⬃1100°C) can increase boundary Cr and Mo levels by ⬃10 wt% over that in the matrix. Isolated measurements on stabilized stainless steels indicate that Nb and Ti may also be enriched by a few percent in the solution-annealed condition. This ‘‘presegregation’’ is commonly thought to result from a nonequilibrium, vacancy drag process with the degree of boundary enrichment dependent on the annealing temperature and the cooling rate. Although data are limited, maximum segregation appears to occur at higher annealing temperatures and at intermediate cooling rates. The mechanism of this presegregation is not well understood, as indicated by Simonen and Bruemmer (22), who demonstrated that Cr enrichments cannot be explained by simple solute–vacancy interactions. Large Cr-vacancy binding energies are required to achieve the observed segregation during quenching, which are completely inconsistent with available data and with nonequilibrium segregation during irradiation. The strongest nonequilibrium segregant in stainless alloys is probably B, which cannot be easily detected by ATEM techniques. Quench-induced B segregation has been observed in stainless steels 46 Jones et al. Fig. 5 Example of nonequilibrium grain-boundary segregation of Mo and Cr during air cooling from annealing temperatures. by atom-probe and radiography techniques and in conjunction with Cr and Mo (20,21). It appears to be likely that B cosegregates with Cr and Mo to grain boundaries during cooling and promotes the presegregation commonly detected due to strong binding between B and vacancies and between B and transition metals. Other elements such as C and N may also reach grain boundaries during cooling and influence presegregation through interactions with Cr and Mo. Interfacial enrichment of Cr (and Mo) should impact the local grain-boundary repassivation behavior and may be critical in the IGSCC resistance of stainless steels. Large heat-to-heat and processing variability can produce boundary Cr concentrations ranging from 18% to 30%. Although enhanced IGSCC resistance might be expected by these large increases in Cr due to improved passivation behavior, much depends on the mechanism of cracking. Austenite stability at the boundary will certainly be altered as well as the oxidation characteristics. In addition, B cosegregation with Cr and Mo to interfaces may influence the chemical and mechanical properties of the grain boundary. Research is needed to assess what role (if any) presegregation plays on the IGSCC of cold-worked stainless steels. Interfacial enrichment of Cr (or other elements that strongly oxidize in preference to the base metal) may be detrimental under specific electrochemical conditions where internal oxidation can occur. Recent results have indicated that this mechanism may control the IGSCC of Ni–Cr stainless alloys in high-temperature water at low electrochemical potentials (23,24). Austenitic Stainless Steels V. 47 CRACK INITIATION PROCESSES The factors controlling crack initiation are expected to be similar to those that control stress-corrosion cracking, namely environment, stress, and material microstructure and microchemistry. However, surface features resulting from fabrication or heat treatment contribute to initiation of cracks but are not a factor in crack growth. Examples of features contributing to stress-corrosion crack initiation in stainless steels have been given by Jones (25), Clarke and Gordon (26), and Licina and Giannuzzi (27). Jones showed a surface feature resulting from surface grinding in preparation for welding of type 304 SS that would allow crevice/occluded cell conditions to occur. Clarke and Gordon found that interface fracture between the matrix and titanium carbonitrides and dissolved silicates and sulfides were potential crack initiation sites in type 304 SS tested in hightemperature water. Licina and Giannuzzi also showed several examples of surface irregularities on ground surfaces of type 304 SS, which would act as crevices/ occluded cells and hence crack initiation sites. Therefore, surface grinding can produce defects that may accelerate crack initiation, although crack initiation will eventually occur at inclusions or grain boundaries if the environment/material/ stress conditions are within a cracking regime. There have been a number of studies aimed at measuring crack initiation. Fatigue crack initiation has been studied extensively, with several studies showing cracks emanating from pits (28). Under stress-corrosion conditions, Clarke and Gordon (26) studied the sites at which cracks initiated in type 304 SS in high-temperature water. They observed that fractures of titanium carbonitride– matrix interface and dissolved silicates and sulfides were sites for crack initiation when these sites were on grain boundaries. Stewart et al. (29) have correlated electrochemical transients with intergranular crack initiation of type 304 SS. They performed slow-strain-rate tests in dilute thiosulfate and found numerous electrochemical current transients that the authors correlated with intergranular crack advance prior to the advance of a single dominate crack. The authors estimated that each current transient was associated with crack advance equal to about one grain diameter and that current decay occurred because of crack arrest. Isaacs (30) used an in situ scanning vibrating electrode technique to identify locations of electrochemical potential variation on the surface of type 304 SS stressed in tension to an unknown level in 10-ppm thiosulfate solution at room temperature. The in situ probe was capable of detecting currents emanating from growing stress-corrosion cracks, and it was shown that several cracks initiated and repassivated prior to a dominate crack initiated and continued to propagate. This result is very consistent with those of Stewart et al. and those obtained with acoustic emission (25). Locci et al. (31) measured the elongation of samples loaded uniaxially in a simple beam apparatus used for creep tests. They measured the crack initiation behavior of ferritic stainless steels in chloride solutions using 48 Jones et al. this apparatus and found the elongation versus time curve correlated with the growth of corrosion trenches to sharp cracks. The initiation period correlated with the transition from flat-bottomed corrosion trenches to 10–20-µm-long sharp cracks emanating from the corrosion trenches in a time period of 10–50 min for a 25% Cr ferritic stainless steel loaded to 90% of its yield strength in boiling 42% LiCl at a potential of ⫺460 mVSCE. VI. FIELD EXPERIENCE Whereas the relative resistance of stainless steels to pitting, crevice corrosion, and SCC are well defined in the laboratory, their applications in the field are far less predictable. This unpredictability is most often the result of poor knowledge of the actual field conditions such as the stress state, the process constituents (i.e., chlorides, oxidizers, pH, etc.), fluctuations in temperature, and process upset conditions. As a result of these inadequacies of knowledge, there are frequent failures of stainless steels in service. Therefore, in order to minimize the potential for in-service failures, it is useful to examine the field performance of the various stainless-steel families, considering both their successes and failures. It is largely by understanding the failures that a better appreciation of the limits of these alloys can be achieved. A. Austenitic Stainless Steels The austenitic stainless steels are the mainstay of most industries. Although Types 304 and 316 are the most commonly used austenitic stainless steels, there are a myriad of standard and specialty austenitics that are used for various applications. The application of these alloys are too numerous to mention here, however, several examples are given to illustrate their use. The chemical process industry uses a large amount of various austenitic stainless steels and, therefore, has accumulated a breadth of experience with these alloys. Figure 6 shows the SCC experience over 10 years with Types 304 and 316 in various plant environments (32). Note that not just the process side can cause SCC but also exposure to nonprocess streams such as steam and cooling water can produce SCC of these alloys. In fact, one of the most significant problems with SCC of austenitic stainless steels for more than 40 years has been the external SCC of austenitic stainless steels under thermal insulation as a result of either chlorides leaching from the insulation or entrapment of chlorides due to exposure to seawater or other chloride-containing waters at various times. The nuclear power industry likewise has had problems with SCC of Type 304 piping and has expended considerable effort to quantify and model the problem. It has been found that among other factors, the solution conductivity has a profound effect on the crack penetration rate (Fig. 7); that is, the lower the con- Austenitic Stainless Steels 49 Fig. 6 Stress-corrosion experience for Type 304 and 316 in various plant environments. ductivity of the solution, the lower the crack penetration rate (33). Factors such as oxygen content, chloride content, and temperature significantly affect solution conductivity and are also difficult to consistently maintain in a plant environment. The need for greater pitting and crevice corrosion resistance of austenitic stainless steels has led to the development of a new stainless-steel family referred Fig. 7 Effect of solution conductivity on crack penetration rate in Type 304 stainlesssteel piping. 50 Jones et al. Table 2 Critical Crevice Temperature (CCT) for Some 6 Month and Other Common Austenitic Stainless Steels Temperature in 10% FeCl 3 –H 2O (pH 1) UNS No. Grade °F °C S31254 N08366 N08367 S30403 S31603 N08904 254 SMO AL-6X AL–6XN Type 304L Type 316L Alloy 904L 90.5 63.5 90.5 ⬍27.5 27.5 32 32.5 17.5 32.5 ⬍⫺2.5 ⫺2.5 0 to as superaustenitics. These alloys generally have ⬃6% Mo and often contain higher N than the typical austenitics. Both of these elements are known to improve the pitting and crevice corrosion resistance of stainless steels. Table 2 shows the enhanced crevice corrosion resistance of these alloys in ferric chloride solution compared to the standard austenitic stainless steels (34). This same improved performance has been observed in seawater service in many field applications. The austenitic stainless steels have been extensively used for equipment in urea plants. The variety of corrosive environments involved with urea synthesis (ammonia, carbon dioxide, urea, and ammonium carbamate) and the associated pressures and temperatures make this process a very complicated one for the application of alloys. An example of such is the rapid failure of Type 316 and cast CF8M (316 equivalent) to ammonium carbamate when small amounts of ferrite are present in the alloy. From experience, it has been found that limiting the ferrite content of CF8M to less than 2% and increasing the chromium content can provide excellent resistance to corrosion from carbamate. However, contrary to this trend and probably as a result of the increased chromium content, the duplex stainless steels provide even better resistance to carbamate corrosion in spite of their high ferrite content. B. Ferritic Stainless Steels The ferritic steels are the simplest of the stainless-steel family of alloys because they are principally iron–chromium alloys. They are widely used for applications that require resistance to atmospheric corrosion and are highly resistant to chloride stress-corrosion cracking. They also provide good oxidation resistance at moderate high temperature. The newer ferritic grades that contain relatively high Austenitic Stainless Steels 51 Cr and Mo additions have found even wider use in the chemical process industry in heat exchangers and other equipment. However, one of the major disadvantages of these alloys is their low strength and non-heat-treatability. Even though the ferritic stainless steels are resistant to SCC from chlorides, they are quite susceptible to pitting and crevice corrosion in environments containing chlorides, so this family of alloys more often fails from localized corrosion rather than SCC. C. Duplex Stainless Steels The duplex stainless steels have gained wide use in many industries over the last 20 years. They not only have generally better corrosion resistance than the austenitics because of the higher Cr content but also have demonstrated superior SCC. Moreover, they have almost twice the yield strength of the standard austenitics. Although they often have greater resistance to SCC than the austenitics, they are certainly not immune. Two recent failures illustrate this limitation. A North Sea offshore separator vessel constructed from 22 Cr duplex stainless steel failed by external SCC after it was determined that seawater was soaking through the insulation onto the hot vessel such that high chlorides then became concentrated on the surface. The vessel operated at 175°F (35). A superduplex stainless-steel manifold was installed in the ocean on a subsea wellhead and exposed to the cathodic protection system which generates hydrogen. Failure occurred quickly as a result of hydrogen embrittlement and plastic strain in critical areas of the manifold connector (36). Duplex stainless steels have been successfully used in urea plants and in the mining industry for portions of pressure acid-leaching plants. They have also been successfully used for pipelines in the oil and gas industry to carry fluids that contain CO2. D. Martensitic Stainless Steels The martensite stainless steels are widely used because of their combined corrosion resistance and the ability to heat treat them to relatively high strength. However, the high strength also makes these alloys prone to hydrogen stress cracking and some forms of SCC. The absence of Ni in these alloys provides essential immunity to chloride SCC from which the austenitic stainless steels often suffer. Type 410 is one of the most commonly used martensitic stainless steels and is used for a variety of applications, including fasteners. At high strength (⬎700 MPa yield), this alloy has suffered from HSC when exposed to seawater and environments containing H2S. Type 420, a higher-carbon-content version of Type 410, is used successfully in the oil and gas industry for tubing to resist 52 Jones et al. corrosion from CO2; however, it is limited to well environments containing little or no H2S. Alloys CA6NM and F6NM have been used for many years for pumps compressors, valves, and so forth and were the precursor to the new family of supermartensitic stainless steels that contain 2–5% Ni and 1–3% Mo. These alloys are considered to have superior pitting resistance in chloride-containing environments compared to the standard martensitic stainless steel. Another martensitic stainless steel that is frequently used is the freemachining grade of Type 410, Type 416, which has reduced pitting and SCC resistance compared to Type 410 due to the large concentration of nonmetallic inclusions. Type 440 stainless steel can achieve very high hardnesses (⬃HRC50), thus providing good wear resistance. This alloy is often used for surgical instruments and cutlery, for which the requirement for corrosion resistance is minimal. However, it is highly susceptible to HSC in any environment in which hydrogen ions are present. REFERENCES 1. ASTM Volume 3.02, Metals Test Methods and Analytical Procedures. West Conshohocken, PA: American Society for Testing Materials, 1998. 2. AJ Sedriks. Corrosion of Stainless Steels, 2nd ed. New York: Wiley–Interscience, 1996. 3. G Cragnolino. Cracking of austenitic alloys. Int Metals Rev 3:85–135, 1979. 4. G Cragnolino, DD Macdonald. Intergranular SCC of austenitic stainless steel at temperatures below 100°C—A review. Corrosion 38:406–424, 1982. 5. DR McIntyre. Experience Survey—SCC of Austenitic Stainless Steels in Water. St. Louis, MO: Materials Technology Institute, 1987. 6. MA Streicher. Analysis of crevice corrosion data from two seawater exposure tests on stainless steels. Mater Perform 22:37–50, 1983. 7. GS Was, SM Bruemmer eds. Grain Boundary Chemistry and Intergranular Fracture, Materials Science Forum, 1989, p. 46. 8. PL Andresen, FP Ford, Mater Sci Eng A103:167, 1988. 9. SM Bruemmer. Corrosion 98, 1998, Paper 138. 10. SM Bruemmer, BW Arey, LA Charlot, Corrosion. 48(1):42, 1992. 11. SM Bruemmer, BW Arey, LA Charlot. In: RE Gold, EP Simonen, eds. 6th Int. Symp. Environmental Degradation of Materials in Nuclear Power Systems—Water Reactors. The Minerals, Metals & Materials Society, 1993, p. 277. 12. SM Bruemmer, LA Charlot, EP Simonen. In: EP Simonen, RE Gold, DE Cubicciotti. 5th Int. Symp. Environmental Degradation of Materials in Nuclear Power Systems— Water Reactors. American Nuclear Society, 1992, p. 821. 13. RH Jones and SM Bruemmer. In: RP Gangloff, MB Ives. Proc. Environment-In- Austenitic Stainless Steels 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 53 duced Cracking of Metals, NACE-10, National Association of Corrosion Engineers, 1990, p. 287. PL Andresen, CL Briant. In: GJ Theus, JR Weeks. Proc. 3rd Int. Sym. Environmental Degradation of Materials in Nuclear Power Systems—Water Reactors. The Metallurgical Society, 1988, p. 371. S Tahtinen, H Hanninen, and M Trolle, ibid 5, pp. 265. TM Angeliu, et al. In: SM Bruemmer, AR McIlree. 8th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems—Water Reactors. American Nuclear Society, 1997, p. 649. P Doig and PEJ Flewitt. Metall Trans A 18A:399, 1987. J Walmsley, P Spellward, S Fisher, A Jenssen. In: SM Bruemmer, AR McIlree, RE Gold. 7th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems—Water Reactors. National Association of Corrosion Engineers, 1996, p. 985. AW James, CM Shepherd. Mater Sci Technol 5:33, 1989. L Karlsson, et al. Acta Metall 36(1):1, 1988. S Dumbill, RM Boothby, TM Williams. Mater Sci Technol 1991, 7:385. EP Simonen, SM Bruemmer. ibid 19, pp. 751. PM Scott, M Le Calvar. Corrosion–Deformation Interactions, EUROCORR 96. The Institute of Materials, 1997, p. 384. SM Bruemmer, LE Thomas, J Daret, PM Scott. Corrosion, 1998. RH Jones. In Proceedings: Workshop on Initiation of Stress Corrosion Cracking Under LWR Conditions. EPRI Report NP 5828. Electric Power Research Institute, Palo Alto, CA, 1988, p. 4-1. WL Clarke, GM Gordon. ibid, p. 9-1. GL Licina, AJ Giannuzzi, ibid, p. 13-1. L Hagn. Mater Sci Eng A 103(1):193, 1988. J Stewart, PM Scott, DE Williams. CORROSION/88, Houston, TX, 1988, Paper 285. HS Isaacs. Initiation of stress corrosion cracking of sensitized Type 304 stainless steel. EPRI report RP1167-8. Electric Power Research Institute, Palo Alto, CA. 1985. IE Locci, HK Kwon, RF Heheman, AR Troiano. Corrosion 43(8):465, 1987. M Nakahara. Preventing stress corrosion cracking of austenitic stainless steels in chemical plants. Nickel Development Institute, 1992. F Ford, PL Andresen. Unresolved modeling issues and their effects on quantitative predictions of environmental cracking. Corrosion/89. National Association of Corrosion Engineers, 1989. RM Davison, JD Redmond. Practical guide to using 6 mo austenitic stainless steel, Mater Perform 27:39, December 1988. I Oystetun, KA Johannson, OB Anderson. Offshore Technology Conf., Houston, TX, 1993, Paper 7207. TS Taylor, T Pendlington, R Bird. Foinaven superduplex materials cracking investigation. Offshore Technology Conf., Houston, TX, 1999. Paper 10965. 3 Nickel-Based Alloys for Resistance to Aqueous Corrosion Paul Crook Haynes International, Kokomo, Indiana I. ADVANTAGES OF NICKEL AS A CORROSION ALLOY BASE In the world of metallic materials for (aqueous) corrosion resistance, nickel and its alloys fill the wide performance gap between the stainless steels and the exotic materials, such as tantalum. Within this same performance band reside the titanium alloys; however, these have more specific uses. The advantages of many nickel-based alloys relative to the stainless steels include the following: 1. Much higher resistance to stress corrosion cracking 2. Higher uniform corrosion resistance, especially in reducing acids, such as hydrochloric, hydrofluoric, and low to moderate concentrations of sulfuric 3. Higher resistance to localized attack (pitting and crevice corrosion), particularly in the presence of chlorides These advantages stem from three attributes of nickel. First, it is more noble than iron; second, it exhibits a ductile, face-centered-cubic (fcc) structure at all temperatures in its solid form; third, it has a high tolerance for useful solutes (alloying additions), such as chromium and molybdenum. The nickel-based corrosion alloys may be grouped in several ways. For the purpose of this chapter, however, they are characterized in terms of the major elemental constituents. Representative, wrought compositions from each group (or 55 56 Crook family) are given in Table 1. From this table, it is evident that the chief alloying elements used in the corrosion-resistant nickel alloys are chromium, copper, molybdenum, tungsten, iron, and silicon. Of these, copper, molybdenum, and tungsten are used to enhance the nobility of nickel; chromium and silicon are used to enhance passivation; iron is used either to lower the cost of the alloys or to shift positions on alloy phase diagrams, so that deleterious microstructural precipitates are avoided. The issue of precipitation within the nickel-based corrosion alloys is extremely complex. However, some discussion of the subject is warranted here. It is generally accepted that, from a corrosion standpoint, the ideal microstructure is one that consists of a single phase. A primary concern, therefore, during the design of the nickel-based alloys has been to avoid phases other than fcc. This has been accomplished not by restricting alloying additions within the soluble range at room temperature but by ensuring that a (metastable) single-phase microstructure is possible at room temperature through a process of annealing and quenching. In other words, the important factors during design have been the high-temperature solubilities and the kinetics of the intermediate-temperature precipitation reactions. As a result of this approach, it is possible for secondary precipitates to form during subsequent excursions to intermediate temperatures (e.g., in heat-affected zones during welding). These precipitates typically form at the grain boundaries of wrought alloys, as they represent ideal nucleation sites. The extent to which these precipitates affect the corrosion resistance and the mechanical properties is related to their atomic structure, shape, composition, and how they influence the composition of the surrounding fcc solid solution. To minimize these precipitation reactions, minor elements such as carbon and silicon are generally held at low levels in the nickel-based wrought alloys. Silicon is a problem, however, in most nickel-based cast alloys for corrosion service, because it is necessary for fluidity. An added problem with castings is elemental segregation (inhomogeneity). This chapter concerns mostly the performance of the wrought, nickel-based alloys. These are typically electric-arc melted, then refined by argon-oxygen decarburization. In many cases, they are remelted (e.g., by the electroslag process) for further refinement and to optimize the structure of the ingot for further processing. Wrought products such as bars and plates are normally made by hot forging and/or hot rolling. Sheets are generally made from hot-rolled coils and are typically cold finished prior to final annealing, to achieve tight tolerances and to control grain sizes. II. CORROSION-RESISTANT, NICKEL-ALLOY SYSTEMS A. Ni Alloys For many applications, the use of commercially pure nickel is warranted. The most common wrought grade is Nickel 200, the composition of which is given Nominal Chemical Compositions of Representative Corrosion-Resistant Nickel-Based Alloys (wt%) Ni Ni alloys Nickel 200 Ni–Cr alloys Alloy 625 Ni–Cu alloys Alloy 400 Alloy K-500 Ni–Mo alloys Alloy B-2 B-3 alloy Ni–Cr–Mo alloys Alloy C-4 C-22 alloy Alloy C-276 C-2000 alloy Alloy 59 Alloy 686 Ni–Cr–Si alloys D-205 alloy Ni–Fe–Cr alloys Alloy 825 Alloy G-3 G-30 alloy Cr 99.0 min. — Cu Mo 0.25 max. — 9.0 Fe W Mn C — max. — 0.40 max. 0.35 max. 0.35 max. 0.15 — 5.0 max. — — max. — 2.50 max. 0.50 max. 2.00 max. 0.3 max. 2.7 2.00 max. 0.50 max. 1.50 max. 0.25 61.0 21.5 — 66.5 66.5 — — 31.5 29.5 69 65 28 — 2 max. 1 max. — 1.5 1.5 0.2 max. 28.5 3 max. 65 56 57 58.5 Bal. Bal. 16 22 16 23 23 21 — — — 1.6 — — 65 20 2.0 2.5 — 42 44 21.5 22 2.2 2.5 3 7 43 30 1.7 5.5 — — Si 16 — 13 3 16 4 — 16 15.75 — 61 3.7 3 max. 3 5 3 max. 1.5 max. 5 max. — — 0.6 0.08 max. 1 max. 0.08 max. 0.5 max. 0.08 max. 1 max. 0.08 max. 0.5 max. 0.1 max. 0.5 max. 0.08 max. 0.75 max. max. — 0.03 — 30 1.5 max. 19.5 0.5 max. 1 max. 1 max. 1 max. 2.5 1 max. 1.5 max. 0.05max. 0.2 max. 0.015 — max. 0.03 max. — — 0.9 — — — — — — Zr — — — — 0.7 max. 0.01 max. — — 0.01 max. — — 0.01 max. — 0.01 max. 0.5 max. — — 0.01 max. 0.25 0.14 0.01 max. — V — 0.1 max. 1 max. 0.01 max. — — — 0.1 max. 3 max. 0.01 max. 0.5 max. 0.2 max. — — 15 — 0.50 max. 0.50 max. 0.10max. 0.40 max. 0.40 max. 3.6 5 6 Cb ⫹ Ta Ti Al — — — 0.2 max. 0.1max. — — — — — — — — 0.35 max. — 0.35 max. — — — — — — — — — — — — 0.5 max. — 0.8 Ni-Based Alloys and Aqueous Corrosion Table 1 — — — — 57 58 Crook in Table 1. Several other grades, with tighter limits on the residual elements, are also available for specific uses (e.g., where enhanced physical or mechanical properties are needed). A review of the physical and mechanical properties of annealed Nickel 200 in Table 2 reveals that this group of materials is characterized by moderate density, a high melting range, low tensile strength, and moderately high-tensile elongation. Fabrication (forming and welding) of commercially pure nickel presents no difficulties, and Nickel 200 is covered by the ASME Boiler Code. A primary use of commercially pure nickel is hardware for handling caustic soda (sodium hydroxide), as it resists practically all concentrations and temperatures of this compound. For example, Nickel 200 is favored for the construction of evaporators, shell and tube heat exchangers, pumps, crystallizers, valves, and fittings used in the concentration and handling of caustic soda (1). Commercially pure nickel is also suitable for caustic potash (potassium hydroxide) service. With regard to the performance of commercially pure nickel in sulfuric acid, the principal applications of Nickel 200 are at room temperature, in unaerated solutions, or where an organic inhibitor is present (2). In hydrochloric acid, Nickel 200 is useful only at room temperature in air-free solutions, up to a concentration of 10 wt%; corrosion rates at higher temperatures are generally unacceptable (3). B. Ni–Cr Alloys This group of materials may also be considered a subset within the Ni–Cr–Mo family. However, they have significantly lower molybdenum contents, so are not quite as resistant to nonoxidizing media and localized attack (pitting, crevice corrosion, and underdeposit corrosion) as the standard Ni–Cr–Mo alloys. Alloy 625 was developed for use as both an aqueous corrosion-resistant material and as an alloy for use in gaseous, high-temperature environments. One of its attributes is moderately high strength, as compared with the Ni–Cr–Mo alloys, allowing the use of thinner structures. The strength level can be controlled, to some extent, by varying the annealing temperature. The strength of Alloy 625 is a result of the presence of niobium (columbium) in the solid solution. Alloy 625 is also slightly age-hardenable, by virtue of the sluggish precipitation of the intermetallic Ni3Cb, although the alloy is not normally used in this condition. Alloy 625 is very popular in the marine and off-shore industries because of its high resistance to localized attack in seawater, as compared with many stainless steels. Undersea applications include exhaust ducts for Navy utility boats and sheathing for communication cables. Alloy 625 is covered by the ASME Boiler Code. C. Ni–Cu Alloys Copper is very soluble in nickel (in fact, they are mutually soluble in all proportions) and enhances its nobility. The basic Ni–Cu alloys, such as Alloy 400, are Physical and Typical Room-Temperature Mechanical Properties of Representative Corrosion-Resistant Nickel-Based Alloys Density Ni alloys Nickel 200 Ni–Cr alloys Alloy 625 Ni–Cu alloys Alloy 400 Alloy K-500a Ni–Mo alloys Alloy B-2 B-3 Alloy Ni–Cr–Mo alloys Alloy C-4 C-22 alloy Alloy C-276 C-2000 alloy Alloy 59 Alloy 686 Ni–Cr–Si alloys D-205 alloy Ni–Fe–Cr alloys Alloy 825 Alloy G-3 G-30 alloy Ultimate tensile strength 0.2% Offset yield strength Tensile elongation (%) g/cm3 lb/in.3 °C °F 8.89 0.321 1435–1446 2615–2635 379–552 55–80 103–207 15–30 40–55 8.44 0.305 1288–1349 2350–2460 827–1034 120–150 414–655 60–95 30–60 8.83 8.47 0.319 0.306 1299–1349 1316–1349 2370–2460 2400–2460 483–621 896–1138 70–90 130–165 172–345 586–827 25–50 85–120 35–60 20–35 9.22 9.22 0.333 0.333 1370–1418 2500–2585 896–917 862–883 130–133 125–128 400–414 400–421 58–60 58–61 55–61 53–58 8.64 8.69 8.89 8.50 8.60 8.73 0.312 0.314 0.321 0.307 0.311 0.315 1357–1399 1323–1371 2475–2550 2415–2500 1310–1360 1338–1380 2390–2480 2440–2516 765–807 765–800 786–793 752–779 690 min. 722–848 111–117 111–116 114–115 109–113 100 min. 105–123 338–421 359–407 359–365 345–393 340 min. 359–421 49–61 52–59 52–53 50–57 49 min. 52–61 52–63 57–70 59–61 62–68 40 min. 56–71 7.99 0.288 1171–1299 2140–2370 786 (sheet) 114 (sheet) 338 (sheet) 49 (sheet) 57 (sheet) 8.14 8.30 8.22 0.294 0.300 0.297 1371–1399 2500–2550 586–724 689 (plate) 676–689 85–105 100 (plate) 98–100 241–448 310 (plate) 310–352 35–65 45 (plate) 45–51 30–50 58 (plate) 55–65 K-500 tensile data in the annealed ⫹ age-hardened condition. MPa ksi MPa ksi 59 a Melting range Ni-Based Alloys and Aqueous Corrosion Table 2 60 Crook therefore ductile, single-phase materials with excellent resistance to non-oxidizing media, such as dilute hydrochloric acid (at low temperatures), hydrofluoric acid, and dilute sulfuric acid. For higher strength levels, age-hardenable Ni–Cu alloys, such as Alloy K500, are available. These contain aluminum and titanium to encourage the controlled precipitation of submicroscopic particles of the intermetallic Ni3(Ti, Al). As may be deduced from Table 2, Alloy K-500 in the aged condition has about twice the strength of Alloy 400. With regard to applications of the Ni–Cu alloys, these are wide and varied, particularly within the chemical process industries and marine engineering. The chemical process industry uses include pressure vessels (Alloy 400 is covered by the ASME Boiler Code), heat exchangers, valves, and pumps. Marine engineering uses include propellers and pumps because the Ni–Cu alloys possess high resistance to degradation in both seawater and brackish water under highvelocity conditions (4). Typical applications of the age-hardenable Alloy K-500 include pump shafts, blades, and scrapers. D. Ni–Mo Alloys The first Ni–Mo alloy (known as Alloy B) was developed in the 1920s for use in cast form. As such, it was limited in performance by elemental segregation and by precipitates induced by significant carbon and silicon contents. Alloy B contained 28 wt% molybdenum and 5 wt% iron, a combination which fortuitously placed the alloy in a fairly safe phase field (i.e., the precipitation of Ni4Mo was avoided). B-2 and B-3 alloys (Table 1) are the modern wrought equivalents. In Alloy B-2, the carbon, silicon, and iron levels were reduced to enhance corrosion resistance, although the reduction in iron caused the alloy to fall in the α ⫹ β phase field, where α is the fcc phase and β corresponds to the ordered intermetallic compound Ni4Mo. This phase is extremely deleterious because it forms rapidly in the temperature range 550–800°C (especially in cold-worked microstructures) and reduces both ductility and resistance to stress-corrosion cracking. By the deliberate addition of minor elements within specific narrow ranges, the precipitation of Ni4Mo has been avoided in B-3 alloy while maintaining the very high corrosion resistance (5). Instead, the γ phase, Ni3Mo, is the stable precipitate. This ordered intermetallic compound takes considerably longer to form, as it requires more diffusion of molybdenum. By virtue of this slower precipitation reaction the B-3 alloy is much more forgiving of slow cooling, following hot forging, hot rolling, and annealing, and it is much more tolerant of elevated temperature excursions during welding. A study of the physical and mechanical properties of the Ni–Mo alloys (Table 2) reveals that they possess moderately high tensile strengths, high ductili- Ni-Based Alloys and Aqueous Corrosion 61 ties, and high melting ranges. With regard to welding and fabrication, care and knowledge of the precipitation propensities are necessary to achieve good results. Both B-2 and B-3 alloys are covered by the ASME Boiler Code. The applications of the Ni–Mo alloys are very specific and related to pure solutions of nonoxidizing acids (organic and inorganic). B-2 and B-3 alloys, for example, are resistant to pure sulfuric and hydrochloric acids at nearly all concentrations and temperatures, up to the boiling points. One of the most important applications, in recent years, has been hardware for making and handling acetic acid. E. Ni–Cr–Mo Alloys The Ni–Cr–Mo alloys are the multipurpose materials of the chemical process and allied industries. They combine the benefits of molybdenum in nickel with the advantages of passivity due to chromium. Whereas the Ni–Mo alloys are unsuitable in the presence of oxidizing contaminants, such as ferric and cupric ions, and dissolved oxygen, the Ni–Cr–Mo alloys can cope with such conditions. They are also extremely resistant to localized attack (pitting, crevice corrosion, and underdeposit corrosion) in the presence of chlorides. From Table 1, it is evident that the chromium contents of the modern wrought Ni–Cr–Mo alloys range from 16 to 23 wt%, and the molybdenum levels range from 13 to 16 wt%, in some cases augmented by tungsten. Other deliberate additions can include iron, to allow the use of less expensive raw materials or to alter positions within the Ni–Cr–Mo alloy phase field, copper, to enhance resistance to nonoxidizing acids, aluminum (for the control of oxygen), manganese (for the control of sulfur), and elements capable of tying up residual carbon, such as titanium and vanadium. Although all the Ni–Cr–Mo alloys are versatile, they each have specific attributes which are taken into account during alloy selection. The high-chromium alloys (C-22, C-2000, and 59), for example, possess much higher resistance to oxidizing media than the low-chromium alloys. Also, a high combined molybdenum plus tungsten content (C-276 and 686) is beneficial in nonoxidizing acids, as is the presence of copper (C-2000). Alloy C-4 has the highest thermal stability (resistance to sensitization). From Table 2, it is evident that the Ni–Cr–Mo alloys possess high ductility, which is helpful in fabricating components. They are also of moderate strength in the annealed condition, thus limiting the thicknesses required for pressure containment. All but the most recently developed Ni–Cr–Mo alloys are covered by the ASME Boiler Code, and applications are in process for the newest materials. With regard to the general performance of the Ni–Cr–Mo alloys, they are resistant to a wide range of chemicals, even in the presence of chlorides. They 62 Crook are generally suitable for use in sulfuric acid up to moderate temperatures [e.g., up to 100°C in the case of C-2000 alloy (except in the concentration range 75– 85 wt%)] (6). The Ni–Cr–Mo alloys are also suitable for use in hydrochloric acid, although the concentration and temperature limitations are more severe. However, several of these alloys provide high resistance to dilute (less than 5 wt%) hydrochloric acid, in some cases up to the boiling points. The Ni–Cr–Mo alloys are among those materials able to handle hydrofluoric acid. Also, they withstand organic acids and alkaline media, such as sodium hydroxide. With regard to industrial applications of the Ni–Cr–Mo alloys, these are many and varied; however, they include reaction vessels, piping, heat exchangers, tanks, valves, nozzles, and pumps. In the power industry, the Ni–Cr–Mo alloys have become the premier materials for lining flue gas desulfurization ducts. Also, in the oil and gas industry, C-276 is used for pipework in some of the harshest downhole conditions. F. Ni–Cr–Si Alloys D-205 alloy, the composition of which is given in Table 1, is the sole representative of this alloy group. It was developed as an alternate to the high-silicon stainless steels, which also possess outstanding resistance to superoxidizing media, such as concentrated commercial sulfuric acid. The advantages of nickel over iron as a basis for a high-silicon alloy include lower work-hardening rates (which are important in the pressing of heat-exchanger plates) and freedom from sigma formation during elevated temperature excursions (instead, a less deleterious intermetallic, Ni3Si, which can actually be used to age-harden the material, forms in D-205 alloy). In the annealed condition, D-205 alloy is characterized by moderate strength and high ductility. Unlike the other wrought nickel-based alloys described in this chapter, however, it is not suitable for use in the as-welded condition because of a continuous, brittle, silicon-rich eutectic phase which forms in the weld metal. The tensile elongation of D-205 weld metal (applied using the gas tungsten arc process) is only about 3%. Annealing at 1040°C (1900°F) is necessary to provide sufficient ductility to these welds for most industrial uses. Table 2 indicates that D-205 alloy also exhibits a relatively low density, compared with the other nickel-based alloys, and a low solidus temperature. This low solidus limits the annealing temperature range. So far, the predominant use of D-205 alloy has been in the form of plate heat exchangers for the cooling of hot, concentrated sulfuric acid, where it has replaced cast iron (cascade coolers) and anodically protected stainless steel (shell and tube heat exchangers). Ni-Based Alloys and Aqueous Corrosion G. 63 Ni–Fe–Cr Alloys Although grouped together because of their major alloying elements, Alloy 825, Alloy G-3, and G-30 alloy (Table 1) have distinctive characteristics and are suitable for different market segments. Alloys 825 and G-3 were designed as multipurpose materials to fill the gap between the high-nickel stainless steels (which are more prone to stress-corrosion cracking) and the Ni–Cr and Ni–Cr–Mo alloys. Both 825 and G-3 offer good resistance to sulfuric acid, by virtue of their combined molybdenum and copper contents. With its higher molybdenum level, Alloy G-3 is significantly more resistant to pitting and crevice corrosion in the presence of chlorides. G-30 alloy was designed specifically for service in commercial phosphoric acid, which is the primary chemical in the agrichemical industries. For this type of service, a high chromium content was found to be desirable, in addition to copper and molybdenum. As a result of its high chromium content, G-30 alloy is also suitable for strong, oxidizing media, such as nitric-acid-based pickling solutions. A review of their respective physical and mechanical properties (Table 2) reveals that the Ni–Fe–Cr alloys are characterized by moderate strengths, in the annealed condition, and high ductilities. All three alloys chosen as representatives of this category are covered by the ASME Boiler Code. Applications of Alloy 825 include equipment for containing and handling sulfuric acid sludges in petroleum refineries (i.e., tanks, heat exchangers, piping, valves, and pumps) and for petrochemical processes which employ phosphoric acid as a catalyst. The main use of G-30 alloy in the agrichemical industries has been for shell and tube heat exchangers. III. NICKEL ALLOYS IN SULFURIC ACID Sulfuric acid is one of the most important industrial chemicals and pervades not only the chemical process industries but also the mining/metal extraction, metal finishing, and agrichemical industries. The performance of the nickel-based alloys in sulfuric acid is very much dependent on acid concentration, temperature, and the presence of other species in the solution. In Ref. 7, three concentration regimes are defined as follows: 1. Low-concentration acid 2. Intermediate-concentration acid 3. High-concentration acid By compiling and comparing data for many nickel-based alloys and stainless steels, it was deduced that the most beneficial elements in low-concentration acid 64 Crook are nickel itself, molybdenum (or tungsten), and chromium. At low concentrations, sulfuric acid is nonoxidizing (reducing), and the cathodic reaction is hydrogen evolution. The intermediate-concentration range is defined as approximately 20–60 wt%. Here too, the cathodic reaction is believed to be hydrogen evolution; hence, the solutions are nonoxidizing. In this concentration range, the Ni–Mo alloys are outstanding, so it is presumed that molybdenum is one of the most beneficial alloying elements. Reference 7 also infers that copper and silicon are beneficial alloying elements within this concentration range. The high-concentration acid is complex, both in electrochemical terms (the redox potential exhibits a steep increase once the concentration exceeds about 70 wt%) and in terms of the influence of the elements on performance. Certainly, the Ni–Mo alloys continue to resist corrosion at these high concentrations, and silicon can become very beneficial above about 90 wt%. However, the corrosion performance of the Ni–Cu alloys deteriorates markedly as the redox potential increases. A review of the data in Ref. 7 leads to the following conclusions: 1. Of the nickel alloys, those based on the Ni–Mo system are the most resistant to pure sulfuric acid, at concentrations up to about 95 wt%; however, these should not be used in aerated systems or where other oxidizing species are present. 2. The Ni–Cr–Mo and Ni–Fe–Cr systems constitute the next best choice at concentrations up to about 95 wt%; these alloys will also resist sulfuric acid in the presence of oxidizing species. 3. In the low- and intermediate-concentration ranges, the Ni–Cu alloys are also worthy of consideration. 4. At very high concentrations of sulfuric acid, the Ni–Cr–Si alloy may be the most suitable alloy, especially in the presence of impurities. Iso-corrosion diagrams for B-3 and C-2000 alloys in sulfuric acid are presented in Figs. 1a and 1b, respectively. These indicate the very safe (0 to 5 mpy), moderately safe (5–20 mpy), and unsafe (over 20 mpy) concentration/temperature regimes for these alloys in pure sulfuric acid. For a perspective, it should be stated that the Ni–Cr–Mo and Ni–Fe–Cr alloys are within the same performance band as Alloy 20 (stainless steel) at concentrations up to about 60 wt%; however, they are considerably better than Alloy 20 in the approximate range of 60–80 wt%. Relative to 316L stainless steel, the Ni–Cr–Mo and Ni–Fe–Cr possess much higher resistance to sulfuric acid at concentrations up to about 85 wt%. A comparison of the Ni–Cr–Si alloy (D-205) and a high-silicon stainless steel (Fe–17 Cr–20 Ni–5 Si) is shown in Fig. 2. This indicates their relative resistance to commercial-grade sulfuric acid at 130°C, in the concentration range Ni-Based Alloys and Aqueous Corrosion 65 (a) (b) Fig. 1 (a) Iso-corrosion diagram for B-3 alloy in sulfuric acid; (b) iso-corrosion diagram for C-2000 alloy in sulfuric acid. 96–99 wt%. A plot of corrosion rate versus temperature for D-205 alloy in 99 wt% acid is shown in Fig. 3; this indicates an upper temperature limit of 150°C. IV. NICKEL ALLOYS IN HYDROCHLORIC ACID Hydrochloric acid is another extremely important industrial chemical. It is both a feedstock and by-product in the organic chemical industry. It is very corrosive, 66 Fig. 2 Crook Comparative corrosion rates in commercial-grade sulfuric acid at 130°C. and only a few alloy systems are suitable for use in the acid at temperatures above ambient. As a result of its volatility, high concentrations are rarely encountered, except in pressurized systems. In a flask/condenser system, for example, the highest concentration that is stable at the boiling point is 20 wt%. Many nickel-based alloys can be used in hydrochloric acid at temperatures Fig. 3 Corrosion rate versus temperature for D-205 alloy in 99% commercial-grade sulfuric acid. Ni-Based Alloys and Aqueous Corrosion 67 (a) (b) Fig. 4 (a) Iso-corrosion diagram for B-3 alloy in hydrochloric acid; (b) iso-corrosion diagram for C-2000 alloy in hydrochloric acid. at and above ambient. Again, their performance is generally a function of temperature, concentration, and the impurities present. The most useful nickel-based alloys for hydrochloric acid service are those of the Ni–Mo and Ni–Cr–Mo systems. As shown in the iso-corrosion diagram for B-3 alloy (Fig. 4a), the Ni–Mo alloys exhibit low corrosion rates in pure (reagent-grade) hydrochloric acid, in the concentration range 1–20 wt%, even up to the boiling points. As in sulfuric acid, however, the Ni–Mo alloys should not be used in aerated systems or in the presence of other oxidizing species, such as ferric ions. The same restrictions 68 Crook apply to the Ni and Ni–Cu alloys, which are useful in low concentrations of hydrochloric acid at low temperatures. The Ni–Cr–Mo alloys not only exhibit low corrosion rates in hydrochloric acid over reasonably wide concentration and temperature ranges (as shown in the iso-corrosion diagram for C-2000 alloy, in Fig. 4b), but they also are tolerant of oxidizing species. However, there are considerable differences in performance between individual compositions within this alloy system, so attention to existing corrosion data or field testing are warranted. V. NICKEL ALLOYS IN HYDROFLUORIC ACID Although not as widely used as hydrochloric acid, hydrofluoric acid is becoming increasingly important as a feedstock and by-product of the organic chemical industry. With it come serious safety concerns, because it such a hazardous compound, and difficulties of containment, because hydrofluoric acid attacks glass. The corrosion rates for several nickel-based alloys in hydrofluoric acid at 79°C (175°F) in the concentration range 5–48 wt% are given in Table 3. Also presented in Table 3, for comparison, are corrosion data for several stainless steels (i.e., the most widely used austenitic grade, one from the high-molybdenum family, one from the alloy 20 group, and one ferritic–austenitic/duplex stainless Table 3 Corrosion Data for Nickel-Based Alloys and Stainless Steels in Hydrofluoric Acid Corrosion rates (mpy) at 79°C (175°F) Alloy 625 Alloy 400 Alloy B-2 B-3 alloy C-22 alloy Alloy C-276 C-2000 alloy Alloy 825 G-30 alloy 316L stainless steel 904L stainless steel 20Cb-3 alloy Alloy 255 5% HF 10% HF 20% HF 48% HF 44.7 16.6 11.8 12.4 25.0 15.9 10.8 31.8 34.0 3,877.0 457.0 30.0 802.0 349.0 16.6 13.8 16.3 31.1 19.0 21.5 44.1 89.3 11,043.0 780.0 29.0 1,528.0 934.0 22.4 17.8 20.7 51.6 34.7 20.2 40.8 83.6 17,760.0 2,449.0 37.3 4,682.0 1,576.0 40.4 25.5 30.6 27.0 36.8 19.4 127.0 269.0 21,336.0 — 59.7 7,247.0 Ni-Based Alloys and Aqueous Corrosion 69 steel). From these data, it is evident that the Ni–Cu alloys, the Ni–Mo alloys, and certain Ni–Cr–Mo alloys can be used in hydrofluoric acid at this temperature, but with moderate rates of attack. The Ni–Cr and Ni–Fe–Cr alloys appear to be unsuitable at this temperature and within this concentration range, as do the austenitic and ferritic–austenitic stainless steels. At much lower temperatures, the Ni–Cr and Ni–Fe–Cr alloys may be used in dilute hydrofluoric acid, as may the copper-containing stainless steel, 20Cb-3 alloy. The performance of the Ni–Cu and Ni–Mo alloys in hydrofluoric acid is strongly influenced by the presence of oxygen, as it is in sulfuric and hydrochloric acids. Considerable increases in the corrosion rates of Alloy 400 in 38% and 48% hydrofluoric acid (boiling) have been recorded, even in the presence of 500 ppm oxygen (9). The Ni alloys, such as Nickel 200, are not as resistant as Alloy 400 to hydrofluoric acid. They are also more strongly affected by oxygen in this acid. Their use is therefore limited to oxygen-free solutions, at temperatures below about 79°C (175°F) (9). VI. NICKEL ALLOYS IN PHOSPHORIC ACID Two distinctly different types of phosphoric acid are encountered in industry (10). The pure (reagent-grade) acid is made from elemental phosphorus, derived from phosphate rock. This is oxidized, then reacted with water. The acid which pervades the agrichemical industries, on the other hand, is made by reacting phosphate rock with sulfuric acid and contains several impurities, such as hydrofluoric acid, sulfuric acid, silica, and chlorides. The levels of these impurities vary depending on the source of the rock, and different batches of this so-called commercial grade of phosphoric acid can vary considerably in their corrosivity. Generally, the commercial grade of phosphoric acid is more corrosive than the reagent grade, and in the commercial grade, a high chromium content has been shown to be extremely beneficial. Test data in 60 wt% commercial phosphoric acid at 116°C are presented for several nickel-based alloys and stainless steels in Fig. 5. These tests were performed under laboratory conditions in a solution provided by an agrichemical company. From these data, it is evident that a high chromium content is beneficial, whether the alloy is nickel based or a stainless steel. Of the materials tested, G30 alloy (30 wt% chromium) exhibited the lowest corrosion rate, whereas Alloy 28 (27 wt% chromium) was the best of the stainless steels. Several nickel-based alloy systems are suitable for use in pure (reagentgrade) phosphoric acid. These include the Ni–Cr, Ni–Cu, Ni–Mo, Ni–Cr–Mo, and Ni–Fe–Cr alloys. 70 Crook Fig. 5 Corrosion rates for nickel-based alloys and stainless steels in 60% commercial phosphoric acid at 116°C. VII. NICKEL ALLOYS IN NITRIC ACID Nitric is a strong, oxidizing acid. Thus, chromium is an extremely beneficial alloying element in nitric acid solutions, as it readily provides passivation. In general terms, the stainless steels are more resistant to nitric acid than the chromium-bearing nickel-based alloys. However, there are occasions when nickelbased alloys are preferred: 1. In heat exchangers, where the stainless steels might not possess sufficient pitting resistance on the cooling water side 2. In multiple-purpose chemical systems, where a batch process involving nitric acid might be followed by another involving a different acid, such as hydrochloric 3. In acid mixtures, where the second acid induces the degradation of stainless steels From Fig. 6, a chart comprising the corrosion rates for several nickel-based alloys and stainless steels in boiling 65% nitric acid, it is evident that care must be taken in choosing a suitable alloy. For example, Alloy C-276, with a chromium content of only 16 wt%, corrodes rapidly in nitric acid. Depending on the requirements for pitting resistance and the other corrosive species encountered, a highchromium Ni–Cr, Ni–Fe–Cr, or Ni–Cr–Mo might be appropriate. Ni-Based Alloys and Aqueous Corrosion 71 Fig. 6 Corrosion rates for nickel-based alloys and stainless steels in 65% boiling nitric acid. VIII. NICKEL ALLOYS IN CHLORIDES Chlorides, even in small amounts, are among the most damaging of chemicals, because they can induce pitting and crevice corrosion in materials which normally exhibit passive behavior. Once initiated, these forms of attack progress at unpredictable rates, often causing equipment to perforate, with a subsequent spillage. Of course, the most abundant chloride-containing solution is seawater, which is commonly used as a coolant in heat exchangers. At high concentrations and temperatures, chlorides can cause stress-corrosion cracking. Chlorides are particularly damaging to the austenitic stainless steels, which are not as resistant to localized attack (i.e., pitting and crevice corrosion) and stress-corrosion cracking as most chromium-bearing nickel-based alloys. Ferric and cupric chlorides are of particular concern, because ferric and cupric ions can markedly alter the nature of the electrochemical process, leading to much higher potentials, hence higher corrosion rates, in the absence of passivation. Thus, alloys such as those in the Ni–Cu and Ni–Mo systems should not be used when these compounds are present. With regard to the suitability of the various nickel-alloy systems for service in chlorides, those in the Ni–Cr–Mo family are the most appropriate, because they not only possess very high resistance to pitting, crevice corrosion, and stresscorrosion cracking, but also they are passive at high potentials, such as those induced by ferric and cupric ions. Alloys in the Ni–Fe–Cr and Ni–Cr families 72 Crook Fig. 7 Critical pitting temperatures for nickel-based alloys and stainless steels in an oxidizing NaCl/HCl solution. exhibit moderate resistance to localized attack and are therefore suitable under mild conditions. To provide some perspective, pitting and crevice corrosion data for several nickel-based alloys and stainless steels are presented in Figs. 7 and 8. The values represent the lowest temperatures at which pitting (Fig. 7) and crevice corrosion (Fig. 8) occur in a solution of 4 wt% NaCl ⫹ 0.1 wt% Fe2 (SO4)3 ⫹ 0.01 M HCl. From these data, it is evident that the temperature required to induce crevice corrosion is generally lower than that needed for pitting to occur. It is also evident that the Ni–Cr–Mo alloys are by far the best nickel-based alloys, with regard to resistance to localized attack in oxidizing chloride solutions. IX. NICKEL ALLOYS IN HYDROXIDES Sodium hydroxide (caustic soda) is the most widely used alkaline material (1). As with acids, the concentration, temperature, and impurities are the most important factors with regard to alloy performance. At low temperatures, iron and steels can be used for handling sodium hydroxide; however, at elevated temperatures, these are subject to caustic embrittlement, and other materials must be considered. Preeminent among the alloy choices for elevated-temperature caustic soda is commercially pure nickel (Nickel 200), which exhibits corrosion rates of less than 1 mpy in boiling solutions up to a concentration of about 50 wt%. Ni-Based Alloys and Aqueous Corrosion 73 Fig. 8 Critical crevice corrosion temperatures for nickel-based alloys and stainless steels in an oxidizing NaCl/HCl solution. With regard to the performance of other nickel alloy systems in sodium hydroxide, it has been reported that the Ni–Cu alloys are practically as resistant as commercially pure nickel (1). Other nickel-based alloys can be used in sodium hydroxide over fairly wide concentration and temperature ranges, although there is some evidence that the Ni–Cr and Ni–Cr–Mo alloys are susceptible to caustic stress-corrosion cracking at very high temperatures and concentrations. In general, the performance of the nickel-based alloys in potassium hydroxide (caustic potash) mirrors that in caustic soda. However, in ammonium hydroxide, which is also important in industry, commercially pure nickel and the Ni– Cu alloys are not recommended, whereas most other nickel-based alloys resist all concentrations up to the boiling points (1,10). REFERENCES 1. Resistance of Nickel and its Alloys to Corrosion by Caustic Alkalies. Corrosion Engineering Bulletin 2, The International Nickel Company. 2. Resistance of Nickel and High Nickel Alloys to Corrosion by Sulfuric Acid. Corrosion Engineering Bulletin 1, The International Nickel Company. 3. Resistance of Nickel and High Nickel Alloys to Corrosion by Hydrochloric Acid, Hydrogen Chloride and Chlorine. Corrosion Engineering Bulletin 3, The International Nickel Company. 4. MONEL Nickel–Copper Alloys, 3rd ed. The International Nickel Company, 1978. 74 Crook 5. DL Klarstrom. A new Ni–Mo alloy with improved thermal stability. Proceedings of the 12th International Corrosion Congress, 1994. 6. P Crook, ML Caruso. The corrosion resistance of Ni–Mo and Ni–Cr–Mo alloys in sulfuric and hydrochloric acids. Proceedings of the 13th International Corrosion Congress, 1996. 7. N Sridhar. Behavior of high-performance alloys in sulfuric acid. Mater Perform 1988; 27(3):40. 8. P Crook, B Ornberg. A new nickel–chromium–silicon alloy for plate heat exchangers. Proceedings of the 10th European Corrosion Congress, EFC, 1993. 9. Corrosion Resistance of Nickel-Containing Alloys in Hydrofluoric Acid, Hydrogen Fluoride and Fluorine. Corrosion Engineering Bulletin 5, The International Nickel Company. 10. N Sridhar. Behavior of nickel-base alloys in corrosive environments. Metals Handbook, 9th ed., ASM International, 1987, Vol. 13, p. 643. 4 Nickel-Based Alloys for Resistance to High-Temperature Corrosion Mark A. Harper Special Metals Corporation, Huntington, West Virginia George Y. Lai Consultant, Carmel, Indiana I. INTRODUCTION Similar to (aqueous) corrosion resistance alloys, the advantages of the Ni-based alloys become evident when the combination of environment plus temperature become too severe for the stainless steels. For materials at high temperatures (e.g., T ⬎ 1000°F), corrosion problems can be very complex, with relatively small amounts of impurities (e.g., Cl in an O2 environment) causing significant changes in the corrosion behavior of a particular alloy. Depending on the environment, which will dictate the general characteristics of the corrosion reaction, eight different corrosion modes can typically be identified in an industrial process. They are oxidation, sulfidation, carburization (including metal dusting), nitridation, halogen corrosion, ash- and salt-deposit corrosion, molten-salt corrosion, and molten-metal corrosion. Although, the criteria for the formation of an oxide scale is simply governed by the thermodynamics of the system (i.e., enough oxygen must be present such that an oxide scale is thermodynamically stable), the continued growth of the scale and its ability to reform when damage, spalling, and so forth occurs is usually governed by kinetic factors. Most industrial environments contain enough oxygen such that the oxidation reaction participates in the corrosion process and a protective oxide scale is usually relied upon by most high-temperature alloys for protection against the various modes of high-temperature attack. Also, it is logical that the formation and growth of the oxide scale and its ability to protect 75 76 Harper and Lai the underlying alloy are more probable in an oxidizing environment (i.e., high oxygen activity, excess oxygen) than in a reducing environment (i.e., low oxygen activity, no excess oxygen). With the formation and growth of oxide scales occurring more slowly under reducing conditions, these types of environments are usually much more aggressive/corrosive than oxidizing environments. Regarding the corrosion reaction(s) occurring in an industrial process, it is important to note that the oxygen activity present in the system is usually an important influence on the rate of attack. For example, sulfidation attack is controlled by both the sulfur and oxygen activities, with an increase in the oxygen activity causing a decrease in the sulfidation attack, and vice versa. Carburization and nitridation behave in a similar manner. However, halogen corrosion, moltensalt corrosion, and molten-metal corrosion behave differently (i.e., oxidizing environments tend to be more corrosive than reducing environments). Thus, an alloy selection process must take into consideration the nature of the corrosive environment and the mode of corrosion attack. Once these two items have been determined, an adequate corrosion database is required in order for a materials engineer to make an informed decision on the appropriate alloy selection. The nominal compositions of several Ni-based wrought alloys used for high-temperature service are shown in Table 1. A more complete list of alloys can be found in Ref. 1. As shown in this table, the levels of carbon, and sometimes nitrogen, in these alloys are relatively high. This is primarily due to the use of the various metal carbides to provide high-temperature strength. Other elements such as Mo, Nb, and W are used for solid-solution strengthening. Although not discussed in this chapter, the high-temperature strength, creep, and thermal fatigue usually play as important a role as the high-temperature corrosion resistance of the alloy. The remainder of this chapter will discuss the various modes of high-temperature corrosion and a review of the available data that exist for Nibased alloys in the various types of environment. Using this information, a materials engineer can make a more informed decision on the appropriate alloy selection for a given application. II. OXIDATION Oxidation is the most common and important high-temperature corrosion reaction, and many alloys rely on the protective oxide scale to prevent attack of the alloy from sulfidation, carburization, and so forth. When the environment contains corrosive impurities (e.g., sulfur, chlorine, etc.), the principal corrosion mode may no longer be oxidation, even though it may still be part of the overall corrosion reaction. For this reason, this section considers oxidation primarily associated with air and ‘‘clean’’ combustion atmospheres generated by using clean Nominal Composition of Alloys (wt%) Alloy name 214 alloy UNS alloy No. C Cr Ni b Co — Fe 3 — 0.05 16 75 600 alloy S alloy 601 alloy 625 alloy 230 alloy 617 alloy X alloy 263 alloy HR-160 alloy HR-120 alloy RA330 alloy 800HT alloy 556 alloy N06600 — N06601 N06625 N06230 N06617 N06002 — N12160 N08120 N08330 N08811 R30556 0.10 a 0.02 a 0.05 0.10 a 0.10 0.07 0.10 0.06 0.05 0.05 0.05 0.08 0.10 15.5 16 23 21 22 22 22 20 28 25 19 21 22 72 c 67 b Bal 62 b 57 b Bal 47 b 52 b 37 b 37 35 32.5 20 — 2a — 1a 5a 12.5 1.5 20 30 3a — — 18 7 3a 14.1 5a 3.0a 1.5 18 0.7 a 3.5 a 33 b Bal Bal 31 b RA85H alloy S30615 0.20 18.5 14.5 — Bal Mo Others W 4.5 Al, 0.5 Mn , 0.2 Sia , 0.1 Zra , 0.01 Ba , 0.01 Y — — 2.25 Nb⫹Ta, 1 Mn a , 0.75 Si a , 0.5 Cu a 15 0.5 Mn, 0.4 Si, 0.25 Al, 0.015 Ba , 0.002 La 1a — — 1.4 Al, 0.5 Mn, 0.2 Si 9 — 3.7 Nb⫹Ta, 0.5 Mna , 0.5 Sia , 0.4 Ala , 0.4 Tia 2 14 0.5 Mn, 0.4 Si, 0.3 Al, 0.02 La, 0.015 B a 9 — 1 Al 9 0.6 1 Mn a , 1 Si a , 0.008 B a 6 — 2.4 Ti a , 0.6 Al a , 0.6 Mn a , 0.4 Si a , 0.2 Cu a a a 1.0 1.0 2.75 Si, 1 Nb a , 0.5 Mn 2.5 a 2.5 a 0.7 Mn, 0.7 Nb, 0.6 Si, 0.2 N, 0.1 Al, 0.004 B 1.25 Si — — 0.8 Mn, 0.5 Si, 0.4 Cu, 0.4 Al, 0.4 Ti — — 3 2.5 1 Mn, 0.6 Ta, 0.4 Si, 0.2 N, 0.2 Al, 0.02 Zr, 0.02 La 3.4 Si, 1 Al — — — — a Ni-Based Alloys and High-Temperature Corrosion Table 1 Note: 214, 230, HR-160, HR-120, and 556 are trademarks of Haynes International. RA330 and RA85H are trademarks of Rolled Alloys, Inc. 800HT is a trademark of the INCO Family of companies. a Maximum. b As balance. 77 78 Harper and Lai fuels (e.g., natural gas or number one and/or number two fuel oil). These fuels generally contain low levels of contaminants such as sulfur, chlorine, alkali metals, and vanadium. Many oxidation problems result from using an alloy in a temperature region that exceeds its protective capability; that is, for a given alloy in service at excessively high temperatures, significant scaling occurs. In response to this problem, a large database on the oxidation of commercial alloys, ranging from carbon steels to superalloys, is available in the literature. Although not discussed in this section, the comparative oxidation resistance of carbon and low-alloy steels, 9 Cr and 12 Cr steels, 17 Cr and 27 Cr stainless steels and austenitic stainless steels at temperatures between 480°C, and 930°C is summarized in The Making, Shaping and Treating of Steel (2). Also, Eiselstein and Skinner (3) summarized the comparative oxidation resistance of austenitic stainless steels, Fe–Ni–Cr alloys, and Ni-based alloys at 980°C. Nickel-based alloys performed significantly better than stainless steels. Most commercial high-temperature alloys rely on a chromia (Cr 2O3 ) scale for protection at elevated temperatures and these alloys are often called ‘‘chromia formers.’’ However, at temperatures above 1000°C, Cr2O3 exhibits significant volatilization to CrO3 and, thus, diminishes the protective capability of a chromiaforming alloy to the oxidation attack. Several chromia-forming Ni-based alloys have been found to exhibit good oxidation resistance at 982°C (1800°F) and higher. For example, Inco’s Inconel 601 and 625 alloys and Haynes International’s Hastelloy S and X alloys, as well as the Haynes 230 alloy, suffered less than 25 µm after 1008 h in air at 982°C. For the same conditions at 1093°C (2000°F), many alloys (e.g., 600, 230, S, and 617) exhibited less than 50 µm of total attack. Even at 1149°C (2100°F), several alloys (e.g., 600, 230, S, and 617) showed less than 100 µm of oxidation attack. However, at 1204°C (2200°F), the oxidation rate increases significantly for the chromia formers, and only a few alloys exhibited less than 250 µm of attack for 1008 h exposure in air. These results, along with the results of other Ni-based alloys and a few austenitic stainless steels, are shown in Table 2 (4). At this temperature, alloys that form and use an alumina (Al2O3) scale for protection are preferred. For example, the 214 alloy exhibited less than 18 µm of oxidation attack when exposed to air at 1204°C for approximately 1000 h. It is important to evaluate an alloy’s long-term oxidation behavior and potential for breakaway oxidation. Harper et al. (5) conducted long-term oxidation studies on several Ni-based alloys at elevated temperatures for periods of 1 and 2 years. Figure 1 shows the weight change as a function of exposure time for the three alloys (HR-120, 800HT, and RA85H) that were exposed to still air at 982°C for 720 days. This figure demonstrates the importance of long-term testing and provides a good example of breakaway oxidation. As shown in this graph, 982°C Alloy Metal loss (µm) Average metal affected a (µm) Metal loss (µm) Average metal affected a (µm) Metal loss (µm) 214 S 230 617 601 600 X 625 556 800H RA 330 310 304 316 1.8 4.6 6.4 7.9 13.5 8.1 8.6 8.1 9.9 23.9 10.2 8.9 140.7 314.2 5.1 12.4 18.0 33.3 32.0 22.9 23.9 18.3 26.7 45.5 108.5 28.7 205.7 362.0 2.0 11.2 11.4 16.3 30.7 27.9 37.8 83.1 24.6 136.9 20.8 24.6 Consumed Consumed 2.0 32.8 32.3 46.5 67.1 41.4 69.1 121.9 65.3 187.7 170.2 57.4 ⬎689.4 b ⬎1737.4b 3.8 25.7 58.2 27.4 59.9 43.9 114.3 405.4 236.5 191.0 40.6 75.4 Consumed Consumed a b 1204°C 1149°C 1093°C Average metal affected a (µm) 4.1 42.2 87.4 85.1 133.9 72.6 148.1 462.3 295.7 255.0 221.7 112.8 ⬎549.9 b ⬎2667 b Metal loss (µm) Average metal affected a (µm) 5.6 Consumed 115.1 269.5 112.3 129.8 Consumed Consumed Consumed 286.3 95.8 201.9 Consumed Consumed 16.5 ⬎805.2 a 201.4 316.7 191.5 213.9 ⬎899.2 b ⬎1209 b ⬎3810 b 344.2 209.6 260.4 ⬎1725.9 b ⬎3566.2 b Ni-Based Alloys and High-Temperature Corrosion Table 2 Metal Loss and Average Metal Affected for Various Alloys Exposed to Flowing Air for 1008 h and Cycled to Room Temperature Once a Week Average metal affected ⫽ Metal loss ⫹ Average internal penetration. Extrapolated to 1008 h. 79 80 Harper and Lai Fig. 1 Weight change versus time for long-term oxidation in still-air testing of HR-120, RA85H, and 800HT alloys at 982°C for 720 days. (From Ref. 5.) for exposures up to 180 days (4320 h), all three alloys appear to be equivalent from a weight change perspective. However, between 180 and 210 days of exposure, the 800HT sample exhibited an accelerated rate of weight loss, characteristic of breakaway oxidation. The RA85H sample exhibited similar behavior after 360 d of exposure. This breakaway oxidation is a result of the transition from a protective Cr 2O3 scale to a less protective Fe,Ni,Cr-spinel oxide scale (6). Table 3 shows the results of the testing conducted at 982°C for a total exposure period of 720 days. The samples were exposed to still air and were subjected to a thermal cycle to room temperature once every 30 days. Figure 2 provides an easy comparison of the alloys exposed in this test and shows the relative amounts of attack that occur via metal loss (i.e., actual thickness loss of the sample) and internal attack caused by internal oxidation and/or void formation. Table 1 also shows the results of similar testing for 360 days of exposure at 1093, 1149, and 1204°C, and Figs. 3–5 show the corresponding graphs of the metal loss plus internal attack. Similar to the results described above, very few chromia-forming alloys can survive longterm exposure at temperatures approaching 1200°C. In contrast, Table 1 and Fig. 5 show the excellent oxidation resistance of an alumina-forming alloy at these high temperatures. 982°C–720 days Alloy Metal loss (mm) Average metal affected (mm) 214 230 617 601 556 HR-160 HR-120 RA85H 800HT — 0.00 0.00 0.01 0.02 0.06 0.04 0.16 0.53 — 0.15 0.24 0.57 0.39 0.42 0.27 1.36 2.03 1093°C–360 days Metal loss (mm) Average metal affected (mm) 0.05 — 0.14 0.36 0.09 0.83 0.45 1.13 Metal loss (mm) 0.27 — 1.15 0.54 0.74 0.97 2.04 1.30 Average metal affected (mm) 0.28 0.54 0.32 Consumed 0.19 1.11 0.51 1.66 Metal loss (mm) 0.86 0.94 1.85 ⬎6.29 1.49 1.35 2.41 1.79 Average metal affected (mm) 0.02 0.01 — — — — 1204°C–360 days 1149°C–360 days 1.51 2.23 0.70 — 0.42 Consumed 0.78 Consumed 2.64 ⬎6.36 2.95 — 2.62 ⬎6.37 ⬎6.39 ⬎6.35 Ni-Based Alloys and High-Temperature Corrosion Table 3 Metal Loss and Average Metal Affected for Various Alloys Exposed to Still Air and Cycled to Room Temperature Once Every 30 Days 81 82 Harper and Lai Fig. 2 Average metal affected for various alloys exposed to still air at 982°C for 720 days. Samples cooled to room temperature once every 30 days. Fig. 3 Average metal affected for various alloys exposed to still air at 1093°C for 360 days. Samples cooled to room temperature once every 30 days. Ni-Based Alloys and High-Temperature Corrosion 83 Fig. 4 Average metal affected for various alloys exposed to still air at 1149°C for 360 days. Samples cooled to room temperature once every 30 days. Fig. 5 Average metal affected for various alloys exposed to still air at 1204°C for 360 days. Samples cooled to room temperature once every 30 days. 84 Harper and Lai Fig. 6 Approximate rate constants as a function of temperature for various oxides. (From Ref. 7.) The results of the 214 alloys shown in Table 2 and 3 are not surprising because alumina formers are known to be more oxidation resistant than chromia formers. Figure 6 shows a plot of the parabolic rate constants for various oxides (7). As shown in this figure, the growth of Al2O3 is at least two orders of magnitude slower than the growth of Cr2O3. However, it should be noted that some studies have shown that these alloys are prone to the formation of internal, randomly distributed voids. In particular, the mechanically alloyed, oxide-dispersion-strengthened (ODS) alloys have been shown to be more susceptible to this type of void formation than conventional wrought alloys (8–10). In addition to the use of alumina-forming alloys at temperatures approaching 1200°C, these alloys are finding use in applications/processes where contamination of the product is a critical issue. An example of this type of application is the furnace equipment used to process electronic components (e.g., semiconductors, capacitors, etc.), glass, and chinaware. In these processes, the major source of contamination is the oxide spalled off of the furnace components, such as wire mesh belts, baskets, and fixtures. These components are typically made from chromia-forming alloys. However, due to the slower growing and more Ni-Based Alloys and High-Temperature Corrosion 85 adherent oxide scale formed on an alumina-forming alloy, furnace equipment fabricated from these alloys results in a much cleaner environment. In addition, these alloys can be put into service with a preformed Al 2O3 scale on the surface and thus eliminate any contamination caused by the transient oxidation of the alloy. For example, preoxidized alloy 214 baskets and wire mesh belts are used for processing of semiconductor components. In terms of product cleanliness, the 214 components perform significantly better than the Ni–Cr alloy previously used (11). III. SULFIDATION One of the most common high-temperature corrosion modes responsible for plant component failures is sulfidation, with two conditions typically responsible for this type of attack. The first involves attack by a gaseous environment, which can be either reducing with H 2 S or oxidizing with SO2. The second involves salt deposits in oxidizing atmospheres with low concentrations of SO2 (less than 1.0% SO2 ). The first condition will be discussed in this section, and the second type of attack will be discussed in the section on ash- and salt-deposit corrosion. One method of approaching the sulfidation problems associated with the attack by gaseous environments is to segregate the environments into three different types: (1) H 2 –H 2S mixtures or sulfur vapor with oxygen activities below the thermodynamic stability region of Cr 2O3 , thus the sulfides are the stable phases, (2) reducing, mixed-gas environments containing H2, H2O, CO, CO2 , H2S, and so forth, and oxygen activities high enough to form Cr 2O3 , and (3) SO2-bearing environments. A. Sulfur Vapor and H2 –H2S Mixtures A review of the sulfidation of metals and alloys in sulfur-vapor and H 2 –H 2 S environments has been conducted by Mrowec and Przybylski (12), Mrowec (13), and Young (14). Most studies have been conducted in sulfur-vapor environments with sulfur pressures greater than 10⫺3 atm and in H 2 –H2 S environments with sulfur partial pressures less than 10⫺2 atm. These sulfur potentials and the very low oxygen activities in the system resulted in the formation of sulfides. Among the Fe–Cr–, Ni–Cr, and Co–Cr alloy systems studied, a significant difference in performance was not noted; however, increasing the chromium content within a particular alloy system generally improved its sulfidation resistance. One application where the H 2 –H 2S mixture is observed is in the gas stream of hydrotreating units for petroleum refining. Severe corrosion attack of the processing equipment has been reported (15–17). The sulfidation behavior of various alloy systems in H 2 –H 2S mixtures has been described by iso-corrosion rate curves 86 Harper and Lai as a function of H 2 S concentration and temperature, with data on chromium steels, Fe–Cr alloys, and austenitic stainless steels reported by Backensto and Sjoberg (18). Sorrell (16) summarized an extensive set of corrosion data for the H 2 –H 2S mixtures typically found in catalytic re-forming units. These data were generated by laboratory, pilot-plant, and field tests and include inspections of commercial operating equipment. Austenitic stainless steels were found to be most resistant followed by straight Cr stainless steels (12–16 wt% Cr). Lowchromium steels (0–9 wt% Cr) were reported to perform poorly. B. Reducing, Mixed-Gas Environments Reducing, mixed-gas environments typically contain H 2 , H 2O, CO, CO2, H 2S, and other gaseous components (e.g., N 2 ) and usually have oxygen and sulfur activities sufficient to form oxides and sulfides on most high-temperature alloys. Thus, the corrosion reaction in these environments usually involves oxidation and sulfidation. In most cases, an alloy will remain relatively protected during an oxidation period where a protective Cr 2O3 scale prevents attack of the alloy by sulfur in the environment. However, at some point, failure of the Cr 2O3 scale, coupled with the alloys inability to quickly reform this scale, results in breakaway corrosion, which is followed by rapid sulfidation attack. Regarding sulfidation problems in coal gasification and similar environments, a large engineering database on commercial alloys was generated during the 1970s and 1980s under several Metal Properties Council’s (MPC’s) programs. These programs evaluated over 80 commercial alloys and coatings, with results documented in MPC annual reports (19) and a summary report published by Howes (20). Overall, high-nickel alloys (e.g., Alloy 600) were very susceptible to sulfidation attack. The Ni–Ni3S 2 eutectic melts at 635°C, and molten sulfide slags can easily destroy the chromium oxide scale and cause catastrophic sulfidation attack. The most important alloying element for improving the sulfidation resistance of iron-, nickel-, and cobalt-based alloys was identified to be chromium. Also, Nagarajan et al. (21) and Norton et al. (22) have studied the effect of silicon on the sulfidation resistance of various Fe–Cr alloys in simulated coal gasification atmospheres. Nagarajan et al. found that an Fe–18 Cr–2 Si alloy performed significantly better than an Fe–18 Cr–0.5 Si alloy in a 24% H 2 –39 H 2O–18 CO–12 CO2 –5 CH4 –1 H 2S atmosphere at 980°C. In a reducing–sulfidizing atmosphere containing 0.8 vol% H 2S at 450°C, Norton found a sharp decrease in the corrosion kinetics of an Fe–12 Cr alloy as the Si content of the alloy was increased from 0.5 to 4.0 wt%. One alloy that has shown excellent sulfidation resistance, both in laboratory and industrial testing, is the HR-160 alloy. This alloy contains high chromium and silicon, along with cobalt (29 Co–28 Cr–2.75 Si), and has been found to perform significantly better than stainless steels and Fe–Ni–Cr alloys (e.g., Alloy Ni-Based Alloys and High-Temperature Corrosion 87 Table 4 Metal Loss and Average Metal Affected for Alloys Tested in a H 2 –25 Vol% CH 4 –14.8 N 2 –4 CO–0.6 CO2 –0.6 H 2S Atmosphere at 899°C for 500 h Alloy Metal loss (mm) Average metal affected (mm) HR-160 HR-120 556 RA85H 253MA 800H 310 0.14 0.59 0.48 0.14 0.44 0.47 Consumed 0.71 1.29 1.30 ⬎3.18 ⬎3.18 ⬎3.18 ⬎3.18 800H) (23,24). One example of this behavior is shown in Table 4. These data were generated from a 500-h exposure to a H 2 –25 vol% CH 4 –14.8 N 2 –4 CO– 0.6 CO2 –0.6 H 2S atmosphere at 899°C (1650°F) (25). In iron-based alloys, aluminum has been shown to be beneficial (26) and Santorelli et al. (27) have reported the sulfidation behavior of two advanced iron-based alumina formers: MA 956 and U.K. Atomic Energy’s Fecralloy alloy. However, nickel-based alumina formers showed poor sulfidation resistance (28). C. SO2-Bearing Environments Most sulfidation studies of SO2-bearing environments have been conducted using pure SO2 or SO2 –O2 mixtures containing high concentrations of SO2. Primarily, pure metals (e.g., Fe, Ni, and Cr) and binary alloys (e.g., Ni–Cr alloys) have been studied. An extensive study of Ni–Cr alloys with various amounts of Cr was conducted by Vasantasree and Hocking (29), with increased amounts of chromium resulting in decreased rates of sulfidation attack. Reviews of the corrosion behavior of metals and alloys have been published by Kofstad (30). However, very few data have been published for commercial alloys. Sulfur furnaces used for the manufacturing of sulfuric acid are the most typical application where environments containing high levels of SO2 are experienced. In this process, sulfur is used as a feedstock for combustion with excess air at approximately 1150°C. The product gas typically contains about 10–15% SO2, (plus 5–10% O2, balance N 2 ), which is then converted to SO3 for sulfuric acid production. One study that examined this type of environment looked at the behavior of a Type 304 stainless steel and the 556 alloy in an oxidizing environment with and without SO2 present (31). Table 5 summarizes the results of this work and suggests that oxidation resistance may be at least one criterion to be used in selecting an alloy for use in an SO2-bearing environment. 88 Harper and Lai Table 5 Metal Loss and Maximum Depth of Attack for Type 304 Stainless Steel and the 556 Alloy Exposed at 982°C for 550 h in an Oxidizing Environment With and Without SO2 Type 304 SS 556 alloy Test gas Metal loss (mm) Maximum depth of attack (mm) Metal loss (mm) Maximum depth of attack (mm) Ar–5 O2 –5 CO2 Ar–5 O2 –5 CO2 –10 SO2 0.31 ⬎0.61 0.44 ⬎0.61 0.005 0.06 0.056 0.10 The other type of environment where SO2 is usually present, albeit in much lower quantities than discussed earlier, is in the combustion of sulfur-bearing fuels such as oil and coal. The combustion flue gas stream produced in a coalor oil-fired boiler typically contains less than 1% SO2. Again, little work has been published for corrosion occurring in these relatively low level SO2-containing environments. One study by Viswanathan and Spengler (32) found that a Ni–15 Cr alloy suffered more attack in a 0.2% SO2 –bal N 2 atmosphere than pure SO2 at 870°C, with the addition of oxygen significantly reducing the corrosion rate. Studies on commercial alloys in environments containing low levels of SO2, particularly those with no excess oxygen, are needed, as this type of atmosphere is relevant to the localized reducing zones that are frequently developed in some industrial boilers. IV. CARBURIZATION AND METAL DUSTING Carburization attack typically occurs when alloys are exposed to an environment containing CO, CH 4 , or other hydrocarbon gases at elevated temperatures. This type of attack results in the formation of internal carbides, causing an embrittlement of the alloy and an overall degradation of its original mechanical properties. Carburization of commercial alloys occurs in many industrial environments; however, most reported problems are experienced in (1) the pyrolysis furnace tubes used in ethylene production and (2) heat-treating equipment such as furnace retorts, baskets, fans, and other components used for the case hardening of steels by gas carburizing. Calciners and furnace components used in the production of activated carbon and carbon fibers are also susceptible to carburization. One of the most effective alloying elements for improving carburization resistance to Fe–Ni–Cr alloys is silicon (33–36), with increased chromium con- Ni-Based Alloys and High-Temperature Corrosion 89 tents generally being beneficial also. Steel and Engel (37) studied Fe–Ni–Cr alloys with chromium contents between 15 and 35 wt%, and found that chromium has a definite beneficial effect on alloys containing 2–25 wt% nickel. Increased chromium levels were less effective for alloys containing 26–45 wt% nickel, and for alloys containing 46–70 wt% nickel, the chromium additions were slightly detrimental. Nickel is known to reduce the diffusivity of carbon in Fe–15 Cr– Ni alloys (38) and increased nickel contents in Fe–Ni–Cr alloys exhibit improved carburization resistance (37). Figure 7 shows this effect of nickel on the carburization resistance of Fe–Ni–Cr alloys. Also, Grabke et al. (39) showed that Fe–Ni– Cr alloys that had a nickel to iron ratio of 4: 1 exhibited maximum carburization resistance, this being in general agreement with the minimum value for the product of the carbon solubility and diffusivity (40). One other beneficial alloying element is aluminum, because alumina formers (e.g., the 214 and MA 956 alloys) have been found to be more carburization resistant than chromia formers (41,42). An example of this behavior is shown in Table 6, where the mass of carbon absorbed by the 214 alloy is compared to other chromia-forming alloys when exposed to a carburizing gas at 1093°C for 24 h (43). As shown in this table, the 214 alloy absorbed approximately one-third the amount of carbon as the best Fig. 7 Effect of nickel on the carburization resistance of Fe–Ni–Cr alloys. (From Ref. 37.) 90 Harper and Lai Table 6 Carbon Absorption for Various Alloys Exposed to an Ar–5 H 2 –5 CO–5 CH 4 Gas at 1093°C for 24 h Alloy Carbon absorption (mg/cm2 ) 214 600 625 230 X S 304 617 316 800H 330 3.4 9.9 9.9 10.3 10.6 10.6 10.6 11.5 12.0 12.6 12.7 performing chromia-forming alloy. The excellent carburization resistance of this alloy has been attributed to the Al2O3 scale formed on the surface of the alloy, which has been confirmed by Auger analysis (41). Similar results have been reported (42) for the MA956 alloy (see Table 7). At temperatures below 815°C, carburization is usually not a problem for industrial equipment because of relatively slow kinetics. However, a problem know as ‘‘metal dusting’’ can occur in strongly carburizing atmospheres (i.e., activity of carbon ⬎ 1) at temperatures generally between 480°C and 815°C. This type of attack exhibits a catastrophic deterioration of metallic materials, with Table 7 Weight Gain for Various Alloys Exposed to a H 2 –2 CH 4 Gas at 1000°C for 100 h Alloy MA-956 601 800 310 Weight gain (mg/cm2 ) ⬍0.3 10.0 19.0 36.0 Ni-Based Alloys and High-Temperature Corrosion 91 Fe-, Ni-, or Co-based alloys decomposing into a ‘‘dust’’ of metal particles, carbon, and sometimes oxides and carbides. Depending on the alloy, maximum attack is observed at approximately 650°C, with alloys typically showing signs of pitting. Metal dusting has been reported for the following: a 1 Cr–0.5 Mo steel in the waste-heat boiler of a partial oxidation syngas production unit at a chemical plant (44); stainless steels in the waste-heat boiler of a synthesis gas reactor (45) and a plant producing gasoline from coal (46); Alloy 800 in the preheater of a gasifier in a coal gasification plant (47) and in a bypass line in a hydrogen reformer (48); Alloy 600 in the re-former of a natural gas–synthetic fuel conversion plant (49); and for various Fe- and Ni-based alloys in carburizing furnaces (50). The environments are typically enriched in H 2 and CO. The fundamentals of metal dusting related to Fe-based alloys have been explained by Grabke et al. (51–53), and a fracture mechanism for metal disintegration during metal dusting has been proposed by Katsman et al. (54). Regarding Ni-based alloys, Grabke et al. (55) have recently published a study on the metaldusting behavior of several commercial alloys and concluded the following: (1) any metallic material is susceptible to metal dusting if carbon is possible at an activity of carbon greater than unity, (2) the mechanism that applies for iron and low-alloy steels does not apply for nickel and Ni-based alloys, and (3) a protective oxide scale or surface poison is required for protection. Also, Klower et al. (56) have found that Ni-based alloys with chromium contents of at least 25 wt% showed no significant evidence of metal dusting for exposures up to 10,000 h, thus supporting work that has shown that a protective chromia scale retards a metal-dusting attack. Compared to the available data on the carburization resistance of commercial alloys, very little data has been published on their metal-dusting behavior. Clearly, more studies are needed not only to develop a better understanding of the mechanism of a metal-dusting attack on Ni-based alloys but also to obtain engineering data regarding the relative resistance of various alloys in various metal-dusting-prone environments. V. NITRIDATION All alloys are susceptible, to some degree, to nitridation attack in ammonia-bearing atmospheres at elevated temperatures and this type of environment is common in the chemical processing industries when ammonia, nitric acid, melamine, and nylon 6-6 (57,58) are produced. Also, ammonia is widely used in the heat-treating industry as a nitriding gas for the case hardening of steels. Similar to carburization, nitridation usually results in the formation of internal nitrides, thus causing an alloy component to become embrittled. When the 92 Harper and Lai Table 8 Nitrogen Absorption and Depth of Nitride Penetration of Various Alloys Exposed to Ammoniaa at 649°C for 168 h Alloy Alloy base Nitrogen absorption (mg/cm2 ) C-276 230 HR-160 600 625 RA333 601 S 617 214 X 825 800H 556 316 310 304 Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Iron Iron Iron Iron Iron 0.7 0.7 0.8 0.8 0.9 1.0 1.1 1.3 1.3 1.5 1.7 2.5 4.3 4.9 6.9 7.4 9.8 a Depth of nitride penetration (mm) 0.02 0.03 0.01 0.03 0.01 0.03 0.03 0.03 0.03 0.04 0.04 0.06 0.10 0.09 0.19 0.15 0.21 100% NH 3 in the inlet gas and approximately 30% NH 3 in the exhaust gas. exposure temperatures are relatively low (e.g., 650°C), a surface nitride layer typically forms (59), with the kinetics of the nitridation process depending on the usual system parameters (i.e., temperature, ammonia concentration in the gas phase, and alloy composition). In general, austenitic stainless steels have been successfully used for the processing equipment in ammonia-bearing environments (58,60–62). However, when the environment is too severe for the stainless steels, nickel-based alloys are known to be more nitridation resistant than iron-based alloys (61). Barnes and Lai (59) studied the nitridation resistance of a wide range of iron-, nickel-, and cobalt-based alloys, with Table 8–10 showing the amount of nitrogen absorption and depth of internal nitride penetration exhibited by the nickel-based alloys when exposed to ammonia for 168 h at 649°C, 982°C, and 1093°C, respectively. Based on this work, Barnes and Lai constructed a plot showing the amount of nitrogen absorption as a function of Ni-Based Alloys and High-Temperature Corrosion 93 Table 9 Nitrogen Absorption and Depth of Nitride Penetration of Various Alloys Exposed to Ammoniaa at 982°C for 168 h Alloy Alloy base Nitrogen absorption (mg/cm2 ) 214 600 S 601 230 617 HR-160 625 X RA333 800H 825 316 556 304 310 Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Iron Nickel Iron Iron Iron Iron 0.3 0.9 0.9 1.2 1.4 1.5 1.7 2.5 3.2 3.7 4.0 4.3 6.0 6.7 7.3 7.7 a Depth of nitride penetration (mm) 0.04 0.12 0.18 0.17 0.12 0.38 0.18 0.17 0.19 0.42 0.28 0.58 0.52 0.37 ⬎0.58 0.38 100% NH 3 in the inlet gas and less than 5% NH 3 in the exhaust gas. nickel plus cobalt in the alloys tested at 982°C (see Fig. 8). Clearly, the beneficial effect of nickel and cobalt on the nitridation resistance of an alloy is evident from this graph. Nitrogen-containing atmospheres can also be nitriding. For example, when exposed to pure nitrogen at 1093°C for 900 h, Alloy 600 and Alloy 800 showed 1.85 mm and 3.81 mm of attack, respectively (63). Swaminathan and Lukezich (64) observed internal nitridation of nickel-based alloys that were exposed to high-velocity combustion gases generated in a gas turbine. Also, Lai (65) observed internal nitridation in four nickel-based alloys (230, 617, 263, and X) that were tested in a laboratory dynamic oxidation burner rig at 982°C for 1000 h. Figure 9 shows an optical micrograph of the cross section of each alloy, with the nitrides highlighted. Certainly, more studies are needed, given that data for the nitridation attack of commercial alloys in this type of environment are rather limited and the fact that nitrogen-containing atmospheres are being used more and more in the production of sintered powder-metallurgy products and other heat-treating operations. 94 Harper and Lai Table 10 Nitrogen Absorption and Depth of Nitride Penetration of Various Alloys Exposed to Ammoniaa at 1093°C for 168 h Alloy Alloy base Nitrogen absorption (mg/cm2 ) 600 214 S 230 617 HR-160 601 625 316 304 X 556 825 RA333 800H 310 Nickel Nickel Nickel Nickel Nickel Nickel Nickel Nickel Iron Iron Nickel Iron Nickel Nickel Iron Iron 0.2 0.2 1.0 1.5 1.9 2.5 2.6 3.3 3.3 3.5 3.8 4.2 5.2 5.2 5.5 9.5 a Depth of nitride penetration (mm) 0.00 0.02 0.34 0.39 ⬎0.56 0.46 ⬎0.58 ⬎0.56 ⬎0.91 ⬎0.58 ⬎0.58 ⬎0.51 0.58 ⬎0.71 ⬎0.76 ⬎0.79 100% NH 3 in the inlet gas and less than 5% NH 3 in the exhaust gas. Fig. 8 Effect of Ni ⫹ Co content on nitridation resistance of Fe-, Ni-, and Co-based alloys. Alloys were exposed to ammonia gas at 982°C for 168 h (100% NH 3 in inlet gas and ⬍5% NH 3 in exhaust gas). (From Ref. 59.) Ni-Based Alloys and High-Temperature Corrosion 95 Fig. 9 Optical micrographs showing internal oxidation and internal nitridation of four alloys tested in dynamic oxidation burner rig at 980°C for 1000 h with a 30-min thermal cycling. (a) Alloy 230, blocky nitrides (arrows), fine carbide precipitates were due to aging during testing; (b) Alloy 617, blocky nitrides (arrows), large needle nitrides and tiny needle nitrides; (c) Alloy 263, tiny needle nitrides; (d) Alloy X, blocky nitrides. (From Ref. 65.) VI. HALOGEN CORROSION In contrast to the behavior of alloys in an environment containing oxygen at high temperatures, where an oxide scale forms and protects the material, the exposure of metallic materials to halogen gases results in the formation of volatile corrosion products consisting of metal halides. In a gaseous environment with both oxygen and a halogen, corrosion of the alloy will involve a combination of an oxide scale and the volatile halides, with the volatile halides causing a significant increase in the spalling and overall degradation of the oxide scale. With respect to industrial processes, the chlorination process has been used to produce titanium, zirconium, tantalum, niobium, and tungsten (66–68), as well as to extract nickel from iron laterites (69) and detinning (70). The manufacture of TiO2 , SiO2 , and ethylene dichloride (EDC) can also involve chlorine. Also, chlorine-containing environments are typically generated during calcining operations used to produce lanthanum, cerium, and neodymium. Thus, for these operations, the reactor vessels, calciners, and other process equipment require alloys that are resistant to high-temperature chlorination attack. It is also important to note that the use of fuels and/or feedstocks that are contaminated with impurities such as chlorine, sodium, potassium, and zinc can result in the reaction of the 96 Harper and Lai Table 11 Melting Point, Boiling Point, and Temperature at Which Certain Metal Chlorides Reach 10⫺4 atm Vapor Pressure Chloride Melting point (°C) Boiling point (°C) 10⫺4 atm temperature (°C) FeCl 2 FeCl 3 NiCl 2 CoCl 2 CrCl 2 CrCl 3 CrO2Cl 2 MoCl 5 WCl 5 WCl 6 TiCl 2 TiCl 3 TiCl 4 AlCl 3 676 303 1030 740 820 1150 ⫺95 194 240 280 1025 730 ⫺23 193 1026 319 987 1025 1300 945 117 268 — 337 — 750 137 — 536 167 607 587 741 611 — 58 72 11 921 454 ⫺38 76 chlorine with the other impurities to form chloride salts. The corrosion that occurs under these will be discussed in Section VII. Returning to the subject of volatile halides, the melting and boiling points of some relevant metal chlorides, as well as the temperatures at which the vapor pressures of these chlorides reach 10⫺4 atm, are shown in Table 11 (71,72). These data illustrate the relatively high volatility and low melting points of the metal chlorides. Nickel and nickel-based alloys are widely used in chlorine-bearing environments. Kane (73) studied the behavior of several commercial iron- and nickelbased alloys in an Ar–30% Cl2 atmosphere at temperatures between 400°C and 704°C and found that higher nickel contents resulted in lower chloridation attack. Table 12 shows the amount of weight loss experienced by various alloys when exposed to the above-mentioned atmosphere and temperatures. The above-discussed data were the result of testing in a chlorine-bearing environment with no measurable oxygen present. However, in many industrial environments, Cl 2 and O2 are usually present, and under these conditions, the formation of both a condensed oxide and the volatile metal halides may occur. The corrosion behavior of various commercial alloys, especially iron and nickel based, in O2 –Cl 2 environments has been studied by several investigators (74– 78). The most notable results of this work has been the observation that alloy Ni-Based Alloys and High-Temperature Corrosion 97 Table 12 Descaled Weight Loss of Several Alloys After Exposure to an Ar–30 Cl 2 Atmosphere at 400, 500, 600, and 704°C for 500 h Descaled weight loss (mg/cm2 ) Alloy 400°C 500°C 600°C 704°C Ni-201 600 601 625 617 800 310 304 347 0.2 0.02 0.3 0.7 0.6 6 28 108 215 0.3 5 3 7 7 13 370 1100 Consumed 47–101 127–180 85–200 — — 200–270 — — — 97 160 215 180 190 890 820 ⬎1000 Consumed Fig. 10 Comparison of the corrosion behavior for the 214, S, and 800H alloys when exposed to an Ar–20 O2 –0.25 Cl 2 atmosphere for 400 h at temperatures between 700°C and 1000°C. (From Refs. 76 and 77.) 98 Table 13 Metal Loss and Total Depth of Attack for Several Alloys After Exposure to an Ar–20 O2 –0.25 Cl 2 Atmosphere at 700, 800, 850, 900, and 1000°C for 400 h 700°C 800°C 850°C 1000°C 900°C Alloy Metal loss (mm) Total depth (mm) Metal loss (mm) Total depth (mm) Metal loss (mm) Total depth (mm) Metal loss (mm) Total depth (mm) Metal loss (mm) Total depth (mm) 214 600 800H 310 S C-276 0.010 — 0.025 — 0.079 0.033 0.010 — 0.033 — 0.081 0.046 0.018 0.020 0.023 0.036 0.145 0.066 0.061 0.086 0.046 0.053 0.150 0.071 0.018 0.038 0.031 0.031 0.224 0.163 0.066 0.132 0.097 0.061 0.257 0.175 0.023 0.127 0.043 0.086 0.315 0.300 0.150 0.252 0.191 0.152 0.353 0.320 0.013 0.330 0.203 0.191 0.419 0.419 0.051 0.386 0.424 0.246 0.472 0.450 Harper and Lai Ni-Based Alloys and High-Temperature Corrosion 99 additions of molybdenum and tungsten, particularly at high concentrations, are detrimental to chlorination resistance. One investigator (75) has attributed this behavior to the formation of molybdenum and tungsten oxychlorides, both of which have very high vapor pressures. The addition of aluminum has been shown to be beneficial to chlorination resistance, especially at temperatures high enough to ensure that the formation and growth of Al 2O3 prevails on alumina-forming alloys. An example of this behavior is shown in Fig. 10, where the total depth of attack as a function of temperature is plotted for three different alloys exposed to an Ar–20% O2 –0.25% Cl 2 atmosphere between 700°C and 1000°C (77,78). The two chromia-forming alloys (800H and Hastelloy S) experienced an increase in attack with increasing temperature; however, the 214 alumina-forming alloy experienced maximum attack at 900°C, with a sudden decrease in corrosion when the alloy was exposed at 1000°C. This has been attributed to the formation and growth of a protective Al 2O3 scale at the higher temperature, whereas at the lower temperatures, the lower growth kinetics of Al 2O3 prevented a completely protective alumina scale from being established on the surface of the alloy. A summary of the results for all the alloys tested in this work is shown in Table 13. The corrosion of several commercial alloys exposed to gaseous HCl has been studied by several investigators (74,79,80). In general, nickel and nickelbased alloys were shown to be more resistant to chlorination attack than ironbased alloys. In contrast to the corrosion studies involving O2 –Cl 2 atmospheres, molybdenum was found to improve an alloy’s resistance to chloridation in reducing environments containing HCl. Hossain et al. (80) found that Ni–Cr–Mo alloys (e.g., 625 and Hastelloy C-4) performed the best among various nickelbased alloys, including nickel. VII. ASH- AND SALT-DEPOSIT CORROSION In many industrial environments, deposits gather on component surfaces, with a subsequent accelerated corrosion attack observed. In most cases, the deposit contains some type of salt, which can cause a damaging chemical reaction between the salt and protective oxide scale. Also, the breakdown of the oxide scale and attack of the alloy can be quite severe when the salt deposit is liquid. The actual formation of the salts usually occurs in the vapor phase, when sulfur and/or chlorine react with other impurities (e.g., K, Na, V, Zn, Pb, etc.) in the fuel or feedstock during combustion. These salt vapors then deposit on cooler components and, in many cases, contain ash from incombustible mineral matter in the fuels. In gas turbines, a type of high-temperature attack known as hot corrosion is both well known and well understood. Sulfur from the fuel reacts with sodium chloride from ingested air during combustion to form sodium sulfate. The sodium sulfate subsequently deposits on hot-section turbine components (e.g., nozzle 100 Harper and Lai guide vanes and rotor blades), causing a fluxing/breakdown of the protective oxide scale that is present on the alloy. The topic of hot corrosion has been extensively reported (81–85) and the mechanism of salt fluxing has been discussed in detail by Goebel et al. (86) and Rapp (87,88). Burner-rig testing is typically used to simulate gas turbine conditions and thus has been used to generate hot corrosion data for several commercial nickeland cobalt-based alloys (89–92). The most important alloying element has been found to be chromium, with alloys containing less than 15% generally showing poor performance in a hot-corrosion-type environment. The effects of other alloying elements was found to be less well defined. In coal-fired boilers, salt deposits may contain sulfur, sodium, potassium, and chlorine, along with the fly ash from the incombustible mineral matter in the coal. The corrosion of superheater and reheater tubes has been very well characterized, with the corrosion rate of austenitic stainless steels exhibiting a bell-shaped curve with respect to temperature (93,94); that is, the rate of attack increases with temperature to a maximum and then decreases with increasing temperature. The formation of molten alkali metal iron trisulfate [(Na,K) 3Fe (SO4 ) 3 ] has been related to the accelerated and then decelerated corrosion shown by the bell-shaped curve. The testing of coextruded tubes using a high-chromium alloy outer layer [310, Inconel 671, and CR35A [45 Ni–35Cr–Fe)] has shown that these alloys are much more resistant than the austenitic stainless steels such as 304H, 316H, 312H, and 347H (95–97). Results from laboratory testing where various alloys were covered with a synthetic ash (37.5 mol% Na 2SO4 , 37.5 K 2SO4 , 25 Fe 2O3 ) and exposed to a synthetic flue gas (80% N2 –15 CO2 –4 O2 – 1 SO2 ) for 50 h at different temperatures are shown in Fig. 11 (97). Another severe deposit-type corrosion is found in oil-fired boilers that burn low-grade fuels (e.g., Bunker C) containing high levels of vanadium, sulfur, and sodium. Typically referred to oil–ash corrosion, vanadium pentoxide and sodium sulfate are the principal phases responsible for this type of corrosion, with reactions between these two constituents resulting in the formation of low-meltingpoint vanadates (98). For uncooled components in the boiler (e.g., hangers and tube supports), alloys containing high levels of chromium, such as 50 Ni–50 Cr, generally perform much better than stainless steels (99–101). In the superheater sections, coextruded tubes containing an outer layer of high-chromium alloys such as 446, 50 Ni–50 Cr, and CR35A also perform better than austenitic stainless steels (102). Fireside corrosion in waste incinerators also involves ash- and salt-deposit corrosion. Corrosive impurities of Cl, S, Na, K, Zn, Pb, and P are usually present in municipal and/or industrial waste. As a result of these impurities in the feedstock, many complex salts, particularly chlorides and sulfates, are formed, thus causing sulfidation and/or chloridation attack. Sulfidation attack of several nickel-based alloys, including 825, 600, 601, 800H, X, and 690, has been ob- Ni-Based Alloys and High-Temperature Corrosion 101 Fig. 11 Laboratory test results for various alloys coated with a synthetic ash (37.5 mol% Na 2SO4 , 37.5 K 2SO4 , 25 Fe 2O3 ) and exposed to a synthetic flue gas (80% N 2 –15 CO2 – 4 O2 –1 SO2 ) for 50 h at different temperatures. (From Ref. 97.) served in municipal waste incinerators (103,104). Also, Whitlow et al. (105) found that chloride-accelerated corrosion attack was responsible for a severe attack of alloys 188, 316 stainless steel, and 825 in a municipal waste incinerator. The HR-160 alloy was found to perform significantly better than many other commercial alloys (106) at temperatures above 650°C (see Table 14); however, for superheater, reheater, and furnace-wall tube applications where the metal temperatures are less than 650°C, recent work has shown that Ni–Cr–Mo alloys perform the best (107). The corrosion reaction at these lower temperatures is still not completely understood. Krause et al. (108) and Krause (109) looked at the corrosion exhibited by various alloys exposed to superheater and furnace-wall conditions. The results showed that molten-salt-deposit corrosion may be a likely mechanism. Certain calcining operations experience ash- and salt-deposit corrosion. 102 Harper and Lai Table 14 Field Test Results from a Municipal Waste Incinerator Where Uncooled Specimens Were Exposed at 700–760°C for 2 months Alloy HR-160 556 625 214 825 304 Metal loss (mm) Maximum metal affected (mm) 0.00 0.24 0.42 0.50 0.55 0.72 0.05 0.28 0.44 0.55 0.62 0.75 Fig. 12 Results of a field rack test in a NaCl–KCl–BaCl 2 salt bath at 840°C for 1 month. (From Ref. 111.) Chemical feedstocks are commonly contaminated with S, Cl, K, and Na impurities, and salts are formed during the high-temperature calcining process. Recuperators used in industrial furnaces can also suffer sulfidation and/or chloride attack. For example, severe corrosion attack can be experienced by stainless-steel recuperator tubes used in aluminum melting operations. The flue gas stream that exits from an aluminum remelting furnace typically contains Cl, S, K, Na, and other impurities as a result of the flux used in the aluminum melting process. These impurities lead to sulfide and/or chloride salt deposits on the recuperator tubes Ni-Based Alloys and High-Temperature Corrosion 103 Table 15 Results of Laboratory Tests in a NaCl Salt Batha at 840°C for 100 hours Alloy 188 556 601 214 304 316 X 310 800H 625 RA330 617 230 S 600 Total depth of attack b (mm) 0.050 0.060 0.060 0.070 0.080 0.080 0.090 0.100 0.100 0.110 0.110 0.120 0.140 0.160 0.190 a Fresh salt bath used for each test run; air used for cover gas. b Mainly intergranular attack; no metal wastage. and rapid attack of the alloy. In one case, a Type 310 recuperator suffered approximately 3.9 mm of attack after 15 months of service at about 650°C. VIII. MOLTEN-SALT CORROSION Separate from salt-deposit corrosion, molten-salt corrosion is related to the corrosion of containment materials that are in contact with a molten salt. Typical examples are molten-salt pots and heat exchangers containing molten salt. The heattreating industry use molten salts extensively for the heat treatment of metals and alloys, with the furnace equipment and other components in contact with the molten salt typically suffering corrosion problems. Other applications of molten salts include heat-transfer and energy-storage media used in solar energy and nuclear systems, high-temperature batteries, fuel cells, and metallurgical extraction processes. 104 Harper and Lai Given that molten salts are usually good fluxing agents which remove oxide scale from a metal surface, the corrosion reaction proceeds by oxidation of the alloy followed by dissolution of the oxide in the molten salt. For this reason, the presence of oxygen and water vapor can accelerate the rate of molten-salt corrosion. This type of corrosion can also take place via mass transfer due to a thermal gradient in the melt. This mode of attack involves the dissolution of an alloying element at hot spots and deposition of that element at cooler spots, which can subsequently result in fouling and plugging in a circulating system. Chloride salts are commonly used in the heat-treating industry for annealing and normalizing of steels at temperatures between 760°C and 980°C. Neutral salt baths, as they are commonly called, typically consist of one or more of the following chlorides: barium, sodium, and potassium. Compositions of five of the more common neutral salt baths are as follows (110): • • • • • 50 50 20 25 21 NaCl–50 KCl KCl–50 Na2CO3 NaCl–25 KCl–55 BaCl 2 NaCl–75 BaCl 2 NaCl–31 BaCl 2 –48 CaCl 2 Jackson and LaChance (110) conducted an extensive study on molten-salt corrosion of cast Fe–Ni–Cr alloys in a 20 NaCl–25 KCl–55 BaCl 2 salt bath. In this study and as typically found in molten chloride salts, they found that the alloys suffered intergranular attack significantly more than metal loss. Also, the results showed that resistance to the molten salt increased with decreasing chromium content and increased nickel content and that the intergranular attack generally followed grain-boundary carbides. Thus, lowering the carbon content of a given alloy could significantly improve its molten-salt corrosion resistance. Lai et al. (111) conducted field testing on various iron-, nickel-, and cobaltbased wrought alloys in a NaCl–KCl–BaCl 2 salt bath at 840°C for 1 month, with the results shown in Fig. 12. In contrast to the above-discussed results for cast alloys, two of the high-nickel alloys (Alloy 600 and Alloy 601) suffered more corrosion attack than the stainless steels 304 and 310. Results from testing in a NaCl salt bath at 840°C are shown in Table 15 (111,112), and similar to the field test results, Co–Ni–Cr–W and Fe–Ni–Co–Cr alloys performed the best. At lower temperatures, corrosion from molten salts typically decreases. Susskind et al. (113) studied various alloys in a molten NaCl–KCl–MgCl 2 salt bath at temperature between 450°C and 500°C and found many alloys resistant to molten chlorides (see Table 16). Nitrate or nitrate–nitrite salt baths are also used for heat-treating purposes, with typical salt bath temperatures between 160°C and 590°C. Applications also exist for use as a medium for heat transfer or energy storage. Slusser et al. (114) evaluated the molten-salt corrosion of various alloys in a NaNO3 –KNO3 salt bath Ni-Based Alloys and High-Temperature Corrosion 105 Table 16 Corrosion Rates of Alloys in Molten Eutectic NaCl–KCl–MgCl 2 Salt at 450–500°C for 1000 h Alloy 1020 2.25 Cr–1 Mo 304 310 316 347 410 430 446 600 N Molybdenum Tantalum Maximum penetration (mm/year) 0 0.08 ⬍0.01 0 ⬍0.01 0.12 0.03 0.05 ⬍0.01 0.05 0.05 0 0.07 at 675°C for 336 h. In general, nickel-based alloys performed better than ironbased alloys. However, pure nickel exhibited a rapid rate of corrosion attack. The corrosion rates of the various alloys plotted as a function of the nickel content are shown in Fig. 13 (114). Longer-term testing (1920 h) showed corrosion rates similar to the 336-h tests, except for Alloy 800, as shown by a comparison of Fig. 13 and Table 17 (114). Also shown in Table 17 is the result of testing at 700°C, where corrosion rates became much higher. Exposure of metals to molten sodium hydroxide (NaOH) results in the formation of metal oxide, sodium oxide, and hydrogen (115). Nickel shows the best resistance to molten NaOH (116–119), particularly low-carbon nickel such as the Ni 201 alloy (120). The corrosion rates of several nickel-based alloys exposed to molten NaOH at temperatures between 400°C and 680°C were reported by Gregory et al. (119) and are shown in Table 18. Based on these results, molybdenum and silicon were detrimental to the molten NaOH salt corrosion resistance. Also, molybdenum and iron were found to be selectively removed from nickelbased alloys with less than 90% nickel, leading to the formation of internal voids (121). The molten-salt nuclear reactor uses a LiF–BeF 2 base salt containing various amounts of UF 4 , ThF 4 , and ZrF 4 , as a fuel salt, and the reactor coolant is a mixture of NaBF 4 –NaF (122). Thus, the corrosion of alloys in molten fluoride salts has been extensively studied for nuclear applications. The most corrosion- 106 Harper and Lai Fig. 13 Corrosion rates of various alloys as a function of nickel content in molten NaNO3 –KNO3 salt at 675°C. (From Ref. 114.) Table 17 Corrosion Rates of Selected Alloys at 675 and 700°C in a Sodium–Potassium Nitrate–Nitrite Salt Corrosion rate (mm/year) Alloy 214 600 N 601 800 675°C 1920 h 700°C 720 h 0.41 0.25 0.23 0.48 1.85 0.53 0.99 1.22 1.25 6.60 Ni-Based Alloys and High-Temperature Corrosion 107 Table 18 Corrosion Rates of Selected Alloys Obtained from Static Tests in Molten Sodium Hydroxide Corrosion rate (mm/year) Alloy 400°C 500°C 580°C 680°C Ni-201 C D 400 600 301SS 0.023 — 0.018 0.046 0.028 0.043 0.033 2.540 0.056 0.130 0.060 0.080 0.06 (a) 0.25 0.45 0.13 0.26 0.96 — (a) — 1.69 1.03 (a) Severe corrosion. resistant alloy in this environment has proven to be the nickel–based alloy N (123). Kroger (122) reported a corrosion rate for this alloy of less than 0.0025 mm per year at 704°C in the LiF–BeF 2 base salt and approximately 0.015 mm per year at 607°C in the NaBF 4 –NaF coolant salt. A variety of commercial alloys were tested on a molten LiF–19.5 CaF 2 salt at 797°C for 500 h by Misra and Whittenberger (124). This salt was being considered for a heat-storage medium in an advanced solar space power system. The tests were conducted in alumina Table 19 Results of Corrosion Tests in LiF–19.5 CaF 2 at 797°C for 500 h Depth of attack (µm) Alloy Mild Steel 304 310 316 RA330 N S X 600 718 188 General a Grain boundary b — — — — — 15 90 — 90 45 — 155 185 130 165 270 15 — 140 30 120 105 Note: Tests were conducted in alumina crucibles under argon. a Intragranular voids near surface. b Intergranular voids. 108 Harper and Lai crucibles with an argon cover gas and the results are shown in Table 19. Chromium appeared to be detrimental in the nickel-based alloys, but no such effect was seen in the iron-based alloys. Molten carbonates are generally less corrosive than molten chlorides or hydroxides. Coyle et al. (125) evaluated a eutectic sodium–potassium–magnesium chloride salt (33 NaCl–21.5 KCl–45.5 MgCl 2 ), a sodium hydroxide salt, and a eutectic sodium–potassium carbonate salt (58 Na 2CO3 –42 K 2CO3 ) for a possible heat-transfer and energy-storage medium capable of operating at 900°C for a solar power generation system. Both the NaCl–KCl–MgCl 2 and NaOH were too corrosive for many commercial alloys; however, the Na 2CO3 –K 2CO3 showed promise because of its less aggressive nature. Results of testing in this molten Table 20 Corrosion Results of Selected Alloys Tested in Molten Eutectic Sodium– Potassium Carbonate at 900°C for 504 h Alloy Total depth of attack a (mm) X 214 188 556 600 b 600 b N 304 316 230 Nickel 800 b 800 b S 0.12 0.19 0.22 0.26 0.34 0.44 0.51 0.54 0.63 0.77 ⬎0.30 0.25 ⬎0.8 ⬎1.43 Note: N2 –0.1 CO2 –(1–10 O2 ) used for cover gas. a All alloys showed metal loss, except for nickel, which suffered 0.2 mm metal loss and ⬎ 0.11 mm intergranular attack. b Two samples from different suppliers. Ni-Based Alloys and High-Temperature Corrosion 109 carbonate salt at 900°C for 504 h are shown in Table 20. The Ni–Cr–Mo Alloy S was severely corroded; the Ni–Cr–Fe–Mo Alloy X performed the best. However, no systematic trend between alloying elements and performance was noted. IX. MOLTEN-METAL CORROSION Similar to molten-salt corrosion, molten-metal corrosion relates to the corrosion of a containment material in contact with a molten metal and/or alloy. Liquid metals are sometimes used as a heat-transfer medium because of their excellent heat-transfer properties and various metals have been investigated for use as a coolant in nuclear reactors. Other applications of liquid metals exist in the heattreating industry (e.g., molten lead) and power generation (126). The molten metal corrosion behavior of a containment material is usually related to its solubility in the molten metal. Thus, a containment material with a higher solubility in the molten metal generally exhibits a higher corrosion rate. Molten aluminum is very aggressive and iron-, nickel-, and cobalt-based alloys are rapidly attacked by this liquid metal. Molten zinc is less of a problem; however, nickel and nickel-based alloys react readily with molten zinc and are not recommended for use in this environment. 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T Flatley, CW Morris, UK Corrosion ’83, Birmingham, UK: Institution of Corrosion Science & Technology, 1982, p. 71. AH Rudd, JM Tanzosh. In: A Armor et al., eds. Proceedings of the 1st International Conference on Improved Coal-Fired Power Plants, Palo Alto, CA: Electrical Power Research Institute, 1986, pp. 2–145. J Stringer. Metals Handbook, 9th ed. Corrosion. Materials Park, OH: ASM International, 1987, p. 998, Vol. 13. WT Reid. In: JM Beer, ed. External Corrosion and Deposits—Boilers and Gas Turbines. New York: American Elsevier, 1971. DW McDowell Jr, JR Mihalisin. ASME Paper 60-WA-260, 1960. BF Spafford. UK Corrosion ’83. Birmingham, UK: Institution of Corrosion Science & Technology, 1982, p. 67. GL Swales, DM Ward. CORROSION/79. Houston, TX: National Association of Corrosion Engineers, 1979, Paper 126. N Bolt. Proceedings of the 10th International Congress on Metallic Corrosion. New Delhi: Oxford and IBH Publishing, 1988, p. 3593. SK Srivastava, GY Lai, DE Fluck. CORROSION/91. 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Houston, TX: National Association of Corrosion Engineers, 1985, Paper 14. 114 Harper and Lai 112. GY Lai. Haynes International, unpublished results, 1986. 113. H Susskind, FB Hill, L Green, S Kalish, L Kukacka, WE McNulty, E Wirsing. Chem Eng Prog 56(3):57, 1960. 114. JW Slusser, JB Titcomb, MT Heffelfinger, BR Dunbobbin. J Metals 24, July 1985. 115. GP Smith. Corrosion of materials in fused hydroxides. USAEC Report ORNL2048, 1956. 116. CM Craighead, LA Smith, RI Jaffee. Screening tests on metals and alloys in contact with sodium hydroxide at 1000 and 1500°F. USAEC Report BMI-705, 1951. 117. EM Simmons, NE Miller, JH Stang, C Weaver. Corrosion and components studies on systems containing fused NaOH. USAEC Report BMI-1118, 1956. 118. RA Lad, SL Simon. A study of corrosion and mass transfer of nickel by molten sodium hydroxide. Corrosion 10(12):435, 1954. 119. JN Gregory, N Hodge, JVG Iredale. 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Washington, DC: US Government Printing Office, 1952, p. 1. 5 Corrosion of Copper and Its Alloys Andrew James Brock Metals Research Laboratories, Olin Corporation, New Haven, Connecticut I. INTRODUCTION Copper and its alloys are some of the earliest metals known to man. The use of copper was recorded in northern Iraq as early as 8500 b.c. and in Egypt in 7000 b.c. Copper articles found in the Sinai Peninsula have been dated as being made in 4000 b.c. The ease with which copper oxide ores can be reduced to copper, together with the ability of the metal to alloy advantageously with other elements, resulted in copper playing a critical role in the development of civilization. The present wide use of copper results from a combination of good corrosion resistance in a multitude of environments, together with other desirable properties. These include its high electric and thermal conductivities, its mechanical properties, and its ease of being formed. Not the least of its uses has been for decorative purposes where the attractive color of the pure element and of alloys such as the brasses, together with their ease of accepting a polish, have considerable aesthetic appeal. Copper and its alloys find other decorative uses such as for the roofs or domes on churches and other institutional buildings. In this case, slow corrosion of the metal at less than 7.5 mils in every 100 years is accompanied by the formation of characteristic green coatings or verdigris. Perhaps the classic example of this is the verdigris on the Statue of Liberty. In other natural environments, such as freshwater and saline waters and soil, copper and copper alloys have high corrosion resistance. They are extensively used for underground water lines, for plumbing tubes, and for condenser tubes in the power utility industry. Copper is a noble metal in that its electrochemical potential is above that of hydrogen. Accordingly, it does not discharge hydrogen from nonoxidizing 115 116 Brock acids. Thus, in oxygen-free sulfuric acid, copper is essentially immune. In the presence of air, dissolved oxygen promotes dissolution of copper by providing an alternate cathodic reaction, namely O2 ⫹ 4H⫹ ⫹ 4e ⫽ 2H2O This supports the anodic reaction of Cu ⫽ Cu2⫹ ⫹ 2 e In neutral solutions, such as freshwater or seawater, oxygen also participates in the corrosion reaction. However, the anodic reaction results in the formation of cuprous oxide: 2Cu ⫹ H2O ⫽ Cu2O ⫹ 2H⫹ ⫹ 2e This oxide film generally forms according to an approximately parabolic rate law and reaches an environment-specific final thickness which is essentially constant. The oxide film provides resistance to the migration of ions and electrons, hence providing a degree of resistance to further corrosion. Cuprous oxide is a p-type semiconductor. The electronic resistance of such oxides can be decreased by doping with higher-valence elements. This is achieved by alloying copper with elements such as aluminum, nickel, tin, zinc, and iron. The oxide films which form on these alloys are then doped with these elements. These doped oxides, together with the presence of corrosion products of the alloying elements, can significantly further increase the corrosion resistance of the metal. A broad range of copper alloys are in use, the alloying elements providing for a required combination of properties. Some of the more common wrought and cast copper alloys and their UNS numbers are listed in Table 1. II. TYPES OF ATTACK Depending on the environment to which the metal is exposed, copper and its alloys can suffer general corrosion or the various types of localized corrosion which are suffered by other metals. It is appropriate to characterize these types of corrosion prior to describing the behavior of the metal in specific environments. A. General Corrosion General corrosion is the term which describes the uniform attack of a metal. It leads to surface roughening on a microscale. Such attack is free of localized corrosion phenomena, which are described below. High rates of localized attack of copper and its alloys result from contact with oxidizing acids. This property is utilized in the cleaning and etching of the metal by immersion in acid–hydrogen peroxide mixtures, nitric–sulfuric acid mixtures, and oxidizing acid salts. Expo- Corrosion of Copper and Its Alloys 117 Table 1 Types and Alloying Constituents in Selected Wrought and Cast Copper Alloys Alloy Wrought Coppers High coppers Brasses Leaded brasses Tin brasses Phosphor bronzes Aluminum bronzes Silicon bronzes Copper–nickels Cast Coppers High coppers Brasses, leaded brasses, and Manganese bronzes Manganese and leaded Silicon brasses and bronzes Tin bronzes and leaded tin bronzes Nickel–tin bronzes Aluminum bronzes Copper–nickels Nickel–silvers UNS No. Alloying elements C10100–C15760 C16200–C10600 C20500–C28580 C31200–C38590 C40400–C49080 C50100–C52400 C60600–C64400 C64700–C66100 C70000–C79900 ⬎99.3% Cu ⬎96% Cu Zn Zn–Pb Zn–Sn–Pb Sn–P Al–Ni–Fe–Si–Sn Si–Sn Ni–Zn C80100–C81100 C87300–C82800 C83300–C85800 ⬎99% Cu ⬎94% Cu Zn–Sn–Pb C86100–C86800 C87300–C87900 C90200–C94500 C94700–C94900 C90520–C95800 C96200–C96800 C97300–C97800 Zn–Mn–Fe–Pb Zn–Sn–Si Sn–Zn–Pb Ni–Sn–Zn–Pb Al–Fe–Ni Ni–Fe Ni–Zn–Pb–Sn sure in environments such as certain soils can result in similar etched surfaces but only after several years. B. Pitting In certain environments, the corrosion of copper alloys can result in pitting. In some cases, pitting occurs over the entire exposed surface of the alloy; in others, pits are formed at discrete locations. In either case, the formation of pits is undesirable because it can lead to structural weakening of the alloy, or perforation in the case of tubes or vessels. Pitting is not the predominant type of attack seen with copper alloys. Environments which typically promote pitting are specific and these will be discussed in later sections. Typically, these environments include sulfide-polluted seawater, water under conditions of stagnation, and certain potable waters with very specific chemistries. The pits often do not continue to grow with time, but attain a certain depth beyond which further increases in depth do not occur. Alloys most resistant 118 Brock to pitting include the high-copper alloys and tin bronzes. Those least resistant to pitting are the low-zinc brasses and aluminum bronzes, whereas the copper– nickels and tin bronzes have intermediate resistance. C. Crevice Corrosion This is a form of attack which can occur where there is a crevice on the exposed alloy surface. Such a crevice can be formed between two copper alloy surfaces or between the alloy and a nonmetallic surface. The latter includes deposits on the metal surface. In the case of copper-alloy condenser tubes, such deposits could be fragments of shells or vegetation or localized deposits of silt. Cleaning of the condenser tubes is one method of eliminating this form of attack. Classically, two types of crevice corrosion are known. In one type, the area inside the crevice is preferentially attacked because it can become depleted in oxygen with respect to the environment outside. The crevice then serves to isolate the anode reaction of alloy dissolution inside the crevice from the cathodic reduction of oxygen outside the crevice. Attack is then essentially due to differential aeration of the metal inside the crevice and that outside it. In the second type of crevice corrosion, initial corrosion of the alloy within the crevice results in an increase in the concentration of metal ions in the solution. On the alloy surface immediately outside the crevice, the metal ions can readily diffuse away or be swept away by the flow of the environment. The overall results is that the areas outside of the crevice suffer dissolution and the metal inside the crevice is ennobled because of the higher metal ion concentration. This form of attack is due to the formation of a metal ion concentration cell. It is the type of crevice corrosion which is more typical of copper alloys. As with pitting, crevice corrosion is a statistical phenomenon. In waters, it is promoted by high temperatures and by the flow of the environment outside the crevice. The depth of attack is usually less than that seen with pitting. D. Dealloying This is a form of localized attack which occurs in alloys where the alloying constituents differ in activity. It has been observed in copper alloys containing Ni, Al, Mn, or Zn. The attack is due to the selective dissolution of the more active component, leaving behind a structurally weak deposit of the less active component. Among copper alloys, dealloying of brasses is most common. It occurs in alloys with more than 15% zinc, the attack being termed dezincification. In waters, this form of attack is favored by high temperature, water stagnation, the presence of crevices where aeration is restricted, a high ratio of chloride ion to bicarbonate ion, and a relatively high pH (1,2). In single-phase brass alloys, the dealloying occurs over the entire alloy Corrosion of Copper and Its Alloys 119 surface and is known as layer-type dezincification (3). In two-phase alloys, the β phase is attacked preferentially at a much higher rate than that of the α phase (4). The corroded metal then contains discrete regions where penetration has occurred, leaving plugs of porous copper. For this reason, such an attack is termed plug dezincification. Figure 1 shows cross sections through dezincified regions of leaded copper–zinc alloys, C36000, which had been immersed in an acidified copper chloride solution for 7 days. It illustrates in one case a uniform dezincified layer. In the other case, dezincification in a cast alloy has penetrated along the β phase which formed at grain boundaries. This represents an extreme case of plug-type dezincification. The dealloying of copper alloys has been the subject of extensive investigation. Early work suggested that the mechanism was one of dissolution of both active and less active components, with the less active component being redeposited on the metal (5,6). Later work suggests that the mechanism was one of selective dissolution of the reactive element (7,8). This mechanism requires diffusion of the more active component through the alloy. Studies of the dealloying of copper–nickel, copper–manganese, and copper–zinc alloys showed that the dealloying kinetics were highest for Cu–Mn alloys and lowest for Cu–Ni alloys, but they were not consistent with diffusion data (8,9). These results led to the conclusion that dealloying rates are primarily controlled by the difference in the potential of the alloy in solution and the reversible potential of the solute element. Layer-type dezincification can be completely inhibited by the addition of low levels of As, P, or Sb (10,11). Bismuth, a similar element present in bismuth brasses, does not play a role in inhibiting dezincification. The role of As, Sb, and P led to the development of Inhibited Admiralty alloys C44300, C44400, and C44500, containing 28% Zn and 1% Sn, with low levels of As, Sb, and P, respectively. However, these elements do not prevent dealloying of the β phase in high-zinc brasses. Tin decreases the rate of attack on the β phase and is found at the 1% level in both naval brass (C46400) and manganese bronze (C67500) for this reason. E. Stress-Corrosion Cracking Stress-corrosion cracking is attack which can occur due to a combination of stress and corrosion where either acting alone would not lead to the development of cracks. With certain susceptible alloys, stress corrosion can lead to very rapid failure. The crack path can be intergranular or transgranular and is perpendicular to the direction of stress. The alloys most susceptible are typically those also most susceptible to dealloying, such as the high-brass and manganese-containing alloys. The attack can result in brittle fracture of the metal. In brasses, this type of failure has also been called season cracking because of the cracking of brass cartridge cases during hot, rainy seasons in former British colonies. Fig. 1 Cross sections of leaded brass alloys showing dezincification after exposure to acidified copper chloride solution for 7 days: (a) shows fairly uniform layer-type attack and (b) shows attack along β phase in a cast alloy. Magnification ⫽ 100⫻. Corrosion of Copper and Its Alloys 121 The atmospheres which promote this form of attack are especially those containing ammonia together with water or water vapor. The presence of oxygen and carbon dioxide accelerate the rate of failure (12). Other corrodents which have been claimed to promote stress corrosion in brasses are amines (13), citrate and tartrate solutions (14), nitrites, carbonates, phosphates, and alkalis (15), sulfur dioxide (16), and seawater (17). Testing for susceptibility to stress-corrosion cracking is typically conducted in Mattssons solution at pH 7.3. This is a solution of ammonium sulfate, copper sulfate, and water with a copper ion content of 0.05 g ions/L and an ammonium ion content of 1.0 g ions/L (18). In one form of the test, strip samples, 30 mils thick and measuring 6 ⫻ 0.5 in. are bent around a 3/4-in. mandrel to provide a permanent 90° set. The sample is then bent into a U and fixed in a jig so that the legs are 3/4 in. apart. The stressed sample is then immersed in the Mattssons solution. Withdrawal of the sample is practiced at intervals. The degree of springback of the samples is compared against that of a control sample not immersed in the solution. Any decrease in springback over that of the control is then due to stress corrosion. When the springback falls to below 80%, the sample is said to have failed. A more rapid test is to expose similar U-bend samples over 28% ammonia in a closed container. Table 2 shows the response of various alloys to such testing (19). The results indicate that many alloys are immune to cracking in the Mattssons solution. Those containing Zn are susceptible, the susceptibility increasing with increasing Zn content. Alloy C66900, containing Mn and Zn, is particularly susceptible. In moist ammonia, other alloys also show susceptibility to stress corrosion. Of the alloys tested, only alloys C11000, C19400, and C65400 and the copper– nickel alloys are immune. Figure 2 illustrates the typical intergranular paths of stress-corrosion cracks which formed in U-bends of alloy C26000 after only 4 h exposure over 28% ammonia. The many branches to the main cracks are characteristic of this corrosion phenomenon. Another test used as an indicator of whether a component will suffer stress corrosion consists of immersion in a mercurous nitrate solution (20). This test is strictly only a means of indicating the presence of residual stress in an alloy, the failure occurring by mercury embrittlement of grain boundaries. As such, it also reveals the presence of residual stresses in alloys, such as the copper–nickels, which are not susceptible to stress-corrosion cracking. Therefore, this test should be confined to testing parts fabricated from brass. F. Corrosion Fatigue Corrosion fatigue, like stress-corrosion cracking, is a form of attack caused by the combination of mechanical and corrosive actions. Metals subjected to cyclic stress in a corrosive environment may be able to withstand a much reduced level 122 Brock Table 2 Stress-Corrosion Data for Selected Copper Alloys in Moist Ammonia and Mattssons Solution, pH 7.3 Time to failure (h) Alloy C11000 C19400 C22000 C23000 C24000 C26000 C26000 C35300 C35300 C42200 C42200 C42500 C42500 C44300 C44300 C51000 C63800 C65400 C66900 C68800 C70600 C71500 C72200 C75200 C75200 C76200 C77000 a Condition Hard Hard Extra hard Cold rolled Spring Hard Extra hard Extra hard Cold rolled Cold rolled Extra hard Cold rolled Spring Cold rolled Cold rolled Extra hard Extra hard Cold rolled Cold rolled Cold rolled Cold rolled Cold rolled Cold rolled Extra hard Cold rolled Cold rolled Extra hard 50% 50% 40% 40% 10% 40% 25% 25% 40% 50% 50% 10% 50% 50% Composition (wt%) 99.9% Cu 2.4% Fe, 0.13 Zn, 0.04 P 10% Zn 15% Zn 20% Zn 30% Zn 30% Zn 35% Zn, 2% Pb 35% Zn, 2% Pb 12% Zn, 1% Sn, 0.2% P 12% Zn, 1% Sn, 0.2% P 9.3% Zn, 2% Sn, 0.2% P 9.3% Zn, 2% Sn, 0.2% P 28% Zn, 1% Sn, 0.04 As 28% Zn, 1% Sn, 0.04 As 5% Sn, 0.2% P 2.8% Al, 1.8% Si, 0.4% Co 3% Si, 1.6% Sn 15% Zn, 12% Mn 23% Zn, 3.4% Al, 0.4% Co 10% Ni, 1.4% Fe 30% Ni, 0.5% Fe 17% Ni, 0.75% Fe, 0.5% Cr 17% Zn, 18% Ni 17% Zn, 18% Ni 30% Zn, 12% Ni 27% Zn, 18% Ni Mattssons a NF NF NF 34 15 4 4.7 2 2 NF NF NF NF 42 13 NF NF NF 0.6 500 NF NF NF 880 530 29 60 Ammonia NF NF 11 2.4 1.2 1.2 3 17 10 37 29 1.2 16 393 28 NF 37 1 NF NF NF — 422 16 — NF ⫽ no failures. of stress compared to the stress level for the same number of cycles in air (21). Failure under such conditions is termed corrosion fatigue. The process is characterized by cracks in the metal which are perpendicular to the tensile stress. Their rate of propagation is usually faster than that of stress-corrosion cracks and they are generally much straighter and with less crack branching. In contrast to stresscorrosion cracks, there is generally only one crack associated with failure by corrosion fatigue. Crack initiation often occurs at the base of corrosion pits. Figure 3 shows a typical fatigue crack initiated on a pit in the outside of an Admiralty alloy condenser tube. Cyclic stresses resulted from vibration of the tube within Corrosion of Copper and Its Alloys 123 Fig. 2 Cross section showing typical branched intergranular stress-corrosion cracks in alloy C26000 after immersion in a moist ammonia environment for 4 h. Magnification ⫽ 320⫻. the condenser, the corrosive environment being ammoniated condensate. The micrographs reveal that the crack is associated with a small pit on the alloy surface and that the crack path is transgranular. Copper alloys most resistant to corrosion fatigue are those with a high fatigue limit and a high resistance to corrosion in the particular environment. Typical of such alloys are the beryllium–coppers, phosphor and aluminum bronzes, and the copper–nickels. G. Intergranular Corrosion As its name implies, this is a form of corrosion in which attack penetrates along grain boundaries, often to a depth of several grains. The more rapid attack of the grain boundaries is generally the result of a difference in composition between the metal in the grain boundary and that in the bulk. Such differences in composition can result from segregation of impurities at the grain boundaries. As intergranular corrosion proceeds in an alloy, the rate of metal loss can increase with 124 Brock Fig. 3 Cross section through the wall of an alloy C44300 tube showing (a) fatigue crack initiating from small pit at 50⫻ and (b) the transgranular nature of crack at 500⫻. Corrosion of Copper and Its Alloys 125 Fig. 4 Cross section of an alloy C44300 condenser tube revealing integranular corrosion on the water side at 1000⫻. time. This is due to grains being isolated and removed from the alloy bulk. Copper alloys most susceptible to this form of attack include Admiralty metal, aluminum brasses, and silicon bronzes. Figure 4 shows an intergranular attack which occurred on the water side of an Admiralty alloy condenser tube which had been in service for close to 30 years with lake water as coolant. Although the depth of attack shown in the micrograph is only some three grains deep, intergranular attack with removal of alloy grains by the flowing water had resulted in a loss in wall thickness of the tube of up to 20 mils. H. Erosion–Corrosion When copper alloys are exposed to an environment under conditions of flow, such as in water tubes or condenser tubes, the flow rate can be significant in determining the corrosion rate. This is because at high fluid velocities, the shear stresses exerted by the fluid on protective films can lead to their removal. Syrett and Lapel (22) defined a breakaway velocity beyond which the effects of erosion cause the oxide to be partly removed. The rate of corrosion in the oxide-free regions increases and is further accelerated by the galvanic action between the oxide-free and oxide-covered areas. The nature of the attack often takes the form 126 Brock of horseshoe-type pits with undercutting in the direction of flow. At higher flow rates, more of the alloy surface becomes free of oxide and the rate of erosion– corrosion decreases because of a decrease in the galvanic effect. Fluid flow in heat-exchanger tubes and condenser tubes is usually fully developed turbulent flow except at the entrance of the tubes. In these regions, local turbulence can promote erosion–corrosion. Additionally, high local fluid velocities can arise around partial blockages in a tube, promoting erosion–corrosion. Erosion–corrosion can also be associated with a phenomenon known as cavitation. This is essentially the wearing away of protective films and metal by repeated impact of blows resulting from the collapse and formation of vacancies or bubbles within a fluid. Such can result from the application of ultrasonics to a fluid and this is utilized in the ultrasonic cleaning of metal. Here, the formation and collapse of vapor bubbles within the fluid results in a scouring of the metal surface with the removal of soils even from recessed parts of the metal surface. The collapse of the vapor bubbles at the metal surface can result in instantaneous stresses of up to 200 ksi. Instances where erosion is promoted by cavitation in corrosive environments include ship’s propellers and in pumps and piping systems subject to severe vibration. Because cavitation is essentially a mechanical form of degradation, the harder forms of copper alloys, such as the aluminum bronzes, are most resistant to it. I. Galvanic Corrosion When two dissimilar metals are immersed in a solution and there is electrical contact between the two, accelerated corrosion of the more electronegative metal can occur while corrosion on the other is reduced. Such accelerated attack is known as galvanic corrosion. Copper and its alloys are usually noble with respect to other metals with which they come into contact. Thus, copper coupled to aluminum or steel will increase the corrosion of the aluminum and the steel while being cathodically protected itself. In these instances, copper is the cathode in the copper–solution–aluminum and copper–solution–steel cells. The degree of attack on the more active component is greatest at the point of contact with the copper. Additionally, the attack is greater, the greater the ratio of the area of the copper member to the more active member. An example of poor design with respect to galvanic corrosion would be the use of steel rivets in copper plates or steel nails in copper roof flashing. Very rapid deterioration of the steel would be anticipated in such cases. The electrochemical potential of a metal in solution is a function of the solution concentration and type. Such potentials are not usually known. Table 3 (23) can be used as a general guideline for designing so that there is no problem with galvanic corrosion. It lists the electrochemical potentials, with respect to a Corrosion of Copper and Its Alloys 127 Table 3 Electrochemical Potential of Metals in Seawater with respect to Saturated Calomel Electrode Alloy Electrochemical potential (V) Magnesium Zinc Aluminum alloys Mild steel–cast iron Low-alloy steel Aluminum bronze Naval brass–yellow brass Tin Copper Lead–tin solder Admiralty brass–aluminum brass Manganese bronze Silicon bronze Nickel silver 90–10 copper nickel Stainless steel 430 Lead 70–30 copper–nickel Nickel 20 Silver Stainless Steel, 302, 304, 321, 347 Stainless steel, 316, 317 Titanium Platinum Graphite ⫺1.6 to ⫺1.63 ⫺0.99 to ⫺1.2 ⫺0.97 to ⫺0.99 ⫺0.6 to ⫺0.71 ⫺0.58 to ⫺0.63 ⫺0.31 to ⫺0.42 ⫺0.3 to ⫺0.4 ⫺0.3 to ⫺0.33 ⫺0.3 to ⫺0.36 ⫺0.28 to ⫺0.37 ⫺0.28 to ⫺0.35 ⫺0.28 to ⫺0.33 ⫺0.25 to ⫺0.28 ⫺0.24 to ⫺0.29 ⫺0.22 to ⫺0.27 ⫺0.2 to ⫺0.27 ⫺0.19 to ⫺0.24 ⫺0.18 to ⫺0.23 ⫺0.1 to ⫺0.2 ⫺0.1 to ⫺0.13 ⫺0.05 to ⫺0.1 0 to ⫺0.1 ⫹0.05 to ⫹0.04 ⫹0.25 to ⫹0.18 ⫹0.3 to ⫹0.2 Source: Ref. 23. saturated calomel electrode, of various metals in seawater. The relative position of the metals will be similar in most neutral solutions and dilute acids. Copper and its alloys occupy the cathodic end of the table. Their corrosion can be galvanically promoted by contact with graphite or titanium. Galvanic corrosion due to contact between two different copper alloys is usually minimal because of the small differences in their electrochemical potentials. The above is illustrated by the results in Table 4, which shows the effect of coupling to other metals on the corrosion rates of some copper alloys in seawater (24). The rate of galvanic corrosion is greatest when the difference in the electrochemical potentials of the two metals in contact is high. Therefore, the coupling 128 Brock Table 4 Corrosion Rates of Alloys in Coupled and Uncoupled Condition After 2 Years of Immersion in Flowing Seawater Alloy Uncoupled C70600 C71500 C61400 Carbon steel Titanium Coupled C70600 C61400 C70600 Carbon steel C70600 Titanium C71500 C61400 C71500 Carbon steel C71500 Titanium Corrosion rate (mils/year) 1.2 0.8 1.7 1.3 0.08 1 1.7 0.12 31 8.2 0.08 0.7 2.5 0.12 28 4.2 0.08 Source: Ref. 25. of such metals in a corrosive environment should not be practiced unless the coupling is via an insulating member. III. OXIDATION OF COPPER ALLOYS In air or oxygen at high temperatures, copper and high-copper alloys oxidize to form films of cuprous oxide, Cu2O, or cupric oxide CuO, or, in many instances, an outer layer of cupric oxide with an inner layer of cuprous oxide. The rate of oxidation and hence the thickness of the oxide films formed in a given time increases with increasing oxidation temperature. At temperatures below 200°C, the rate of oxidation of copper has been described in terms of an inverse logarithmic law (25) and a logarithmic law (26). The expressions are Corrosion of Copper and Its Alloys 冢 冣 inverse logarithmic law 冣 logarithmic law 1 t ⫽ k1 log ⫹1 m k2 m ⫽ k 3 log 冢 t ⫹1 k4 129 where m is the weight of oxide formed, t is the time, and the other terms are constants. At these temperatures, the thickness of the oxide film is usually of the order of the wavelength of light. Because of this, oxides formed are often translucent and colored because of the property of light interference. The thickness of the oxide films can be estimated from the color of the oxide as illustrated in Table 5 (27). At higher temperatures, the rate of oxidation is usually described in terms of a parabolic rate equation (28) and sometimes in terms of a cubic rate equation (29). The expressions are m 2 ⫽ k 5 t ⫹ k 6 parabolic rate equation m 3 ⫽ k 7t ⫹ k 8 cubic rate equation where the terms have the same meaning as given previously. At temperatures in the range 350–1000°C, thick oxide films are formed which spall off from the metal when the samples are cooled. Table 5 Color of Copper Oxide Films as Function of Thickness Oxide color Dark brown Red brown Dark purple Dark violet Dark blue Pale blue-green Pale silvery green Yellowish green Yellowish green Old gold Orange Red brown Source: Ref. 27. Thickness ˚) (A 380 420 450 480 500 830 880 970 980 1110 1200 1260 130 Brock Thermodynamic considerations would suggest that oxidation in air at 1 atm and temperatures below 300°C would favor the formation of cupric oxide, with cuprous oxide being favored as an increasing fraction of the overall film as the temperature increases above 300°C. Results reported in the literature are generally in accord with this. Thus, Hickman (30), using electron diffraction studies, demonstrated that between 20°C and 300°C, cupric oxide films were formed. At 350°C, films consisting of both cuprous oxide and cupric oxide were identified, and above 450°C, the films consisted entirely of cuprous oxide. Alloying copper with elements such as Fe, Mn, or Ni has only a minor effect on their oxidation behavior (31). Additions of Al, Be, Mg, Si, or Zn can significantly decrease the oxidation rates. This is because oxides of the alloying constituents form beneath the copper oxide layer. At high-alloying contents, such as 15% Zn (32) or 5% Al (33), oxidation at temperatures in excess of some 500°C results in the formation of a complete layer of zinc oxide or alumina adjacent to the surface, which significantly decreases the oxidation rate of the metal. Copper alloys are not used as materials for oxidation resistance at high temperatures. However, the above types of oxides can result when the alloys are heat treated in air at high temperatures. Removal of copper oxides is generally accomplished by immersion in 10–20% sulfuric or hydrochloric acids. Refractory oxides, such as silica or alumina, must be removed by treating with 10–20% sulfuric acid made oxidizing by the addition of 3% hydrogen peroxide. In this case, the oxides are removed by dissolution of the alloy layers under them. The formation of copper oxides can be reduced by annealing in nitrogen atmospheres and can be completely prevented by annealing in nitrogen containing 1–4% hydrogen. Such reducing atmospheres are useful because they do not form explosive mixtures with air. Reducing atmospheres with much higher hydrogen contents can be used to anneal brasses without oxidation of Zn. However, these atmospheres will not prevent the oxidation of reactive elements such as Si, Al, or Mg. These either oxidize to form internal oxides, often preferentially at grain boundaries, or they form uniform layers of the oxide of the alloying elements. Transition from internal to external oxidation is favored by increasing the temperature and increasing the alloying content of the reactive phase. Figure 5 demonstrates this for copper–silicon alloys with increasing silicon content oxidized at 600°C in an atmosphere of nitrogen–4% hydrogen (19). IV. ATMOSPHERIC CORROSION Copper and its alloys have a high resistance to corrosion in the atmosphere. Although copper oxidizes on exposure to the atmosphere, corrosion is prevented by the formation of an adherent layer of corrosion products. In some instances, Corrosion of Copper and Its Alloys 131 Fig. 5 Cross sections of copper silicon alloys oxidized for 16 h at 600°C in nitrogen– 4% hydrogen showing transition from internal oxidation to external film formation as silicon content increases from (a) 0.2% through (b) 0.5% to (c) 2.5% at bottom. Magnification ⫽ 800⫻. this coating takes the form of a green patina. On copper on roofs or statues, this patina provides for an aesthetically pleasing appearance. Vernon and Whitby (33) characterized the composition of the patina formed on copper exposed in England for times ranging up to 300 years. Under most conditions of exposure, they showed that the major constituent was brochantite, a basic copper sulfate. A basic copper chloride, atacamite, was found in the corrosion products formed on copper near the sea. Lesser amounts of malachite, a basic copper carbonate, were also detected. It is evident from work described in the literature that the rate of corrosion of copper in the atmosphere depends on the specific environment, be it urban, marine, or rural, Other factors which play a role are the temperature and humidity, the degree of pollution, and the relative time spent between wet and dry conditions. From measurements of the increase in resistance of copper and copper- 132 Brock alloy wires exposed in an urban atmosphere, Hudson (34) showed that over a time of close to 5 years, copper corroded at an average rate of 0.195 mils/year Slightly lower corrosion rates were observed with silicon and tin bronzes but significantly higher corrosion rates for 80–20 and 70–30 cupronickel alloys, 70– 30 brass, and an aluminum bronze. Thompson et al. (35) used weight-loss measurements to determine the rate of corrosion of sheet and wire samples of 11 different copper alloys in various atmospheres over a time of up to 2 years. Over this time, they determined average corrosion rates of 34–47, 27–37, and 12–17 µin. per year in industrial, marine, and rural atmospheres, respectively. Tracy (36) also summarized the corrosion rates of copper and several copper alloys exposed for 20 years at eight different sites. For most of the alloys, the corrosion rates were within the ranges 2–30, 10–90, and 50–120 µin. per year. Exceptions to the behavior were two brasses, for which the rates of attack were 180–450 mil/year. Holm and Mattsson (37) evaluated sheet and rod samples of a wide range of copper alloys after exposure for times up to 16 years in rural, marine, and urban atmospheres in Sweden. They observed that dark brown coatings consisting essentially of copper oxides together with some copper salts and alloying constituents formed in the first year of exposure. In urban and marine environments, green coatings appeared on copper after 6–7 years and even earlier on phosphor bronzes. However, green coatings were not formed on high-zinc alloys. Such coatings were not formed on any of the alloys exposed at rural sites even after 16 years exposure. Instead, black or brown films were formed. They showed that the average rate of loss of metal by general corrosion after 7 years was much lower than that after only 2 years of exposure but similar to that after exposure for 16 years. The values were 12–20 µin./year in rural atmospheres, 20–35 µin./year in marine atmospheres, and 35–52 µin./year in urban atmospheres. Penetration of the brasses by dezincification occurred at rates significantly higher than those described. It was greatest for two-phase alloys and after 16 years had penetrated to 3.5–5.6 mils. Their results showed that two-phase alloys containing additions of aluminum, tin, or arsenic improved the dezincification resistance. Single-phase brasses containing arsenic additions were shown to have good resistance to dezincification. Costas (38) evaluated a range of copper alloys after exposure for up to 20 years in rural, marine and industrial locations. He observed that at industrial sites, green coatings predominated on samples free of zinc or nickel. Blue or green hues were seen on a few alloys exposed at both marine and rural sites but not to the same degree as at the industrial sites. The average corrosion rates varied from 9 to 90 µin./year and were greatest at the industrial site. Figure 6 shows weight loss–time plots for alloy C26000 exposed at three different sites and illustrates the dependency of the corrosion rate on the site Corrosion of Copper and Its Alloys 133 Fig. 6 Weight loss–time plots for alloy C26000 exposed to the atmosphere at three different sites. conditions (19). The New Haven site is representative of an industrial atmosphere, Daytona Beach represents a marine environment with negligible industrial pollution, and East Alton represents a severe industrial atmosphere. The corrosion rate of the brass is much less at the Daytona Beach site, reflecting the absence of industrial pollutants. Figure 7 shows similar plots for three alloys exposed at Daytona Beach. The corrosion rate of the high-copper alloy, C11000, is significantly greater than those of brass, C26000, and of the copper–nickel alloy, C70600. The plots also demonstrate that, with time, the corrosion rates generally Fig. 7 Weight loss–time plots for selected copper alloys exposed in New Haven. 134 Brock Table 6 Sites Corrosion Rates of Copper Alloys After 2 Years of Exposure at Various Alloy C11000 C22000 C23000 C26000 C42500 C51000 C65400 C68800 C70600 C76200 Corrosion rates (mils/year) Alloying elements (wt%) Daytona Beach New Haven East Alton 99% Cu 10% Zn 15% Zn 30% Zn 9.3% Zn, 1% Sn, 0.2% P 5% Sn, 0.2% P 3% Si, 1.6% Sn 2.8% Al, 1.8% Si, 0.4% Co 10% Ni, 1.4% Fe 3% Zn, 12% Ni 0.1 0.07 0.03 0.03 0.08 0.15 0.28 0.02 0.06 0.02 0.04 0.03 0.03 0.05 0.04 0.05 0.04 0.03 0.03 0.05 0.04 0.06 0.04 0.04 0.06 0.04 0.04 0.03 0.06 0.04 Source: Ref. 19. decrease because the corrosion product layers provide some resistance to further attack. The rate of corrosion at any time is given by the slope of a particular plot at that time. Table 6 shows the corrosion rates so determined for several alloys after 2 years of exposure at the three sites. Stress-corrosion cracking of certain copper alloys can also occur during exposure to the atmosphere. Table 7 shows stress-corrosion data for copper-alloy samples exposed as U-bends at New Haven, Daytona Beach, and Brooklyn, NY. In contrast to the data shown in Table 3 for laboratory testing in moist ammonia and the Mattssons solution, brasses with 15% or less Zn did not fail within the 10-year exposure time. However, failures were observed for alloy C66900 with only 15% Zn but with 12% Mn. The results also demonstrate a sensitivity to the site, with the marine site being least aggressive in promoting stress-corrosion cracking. Under certain conditions, much more rapid atmospheric corrosion of copper alloys can occur. Typically, such conditions include the presence of sulfides such as hydrogen sulfide, which leads to the formation of black, much less protective films of copper sulfide. V. WATER STAINING Very rapid staining of copper alloys can occur under specific conditions. Typically, this occurs on freshly cleaned strip or parts which have not been fully Corrosion of Copper and Its Alloys 135 Table 7 Time to Failure by Stress-Corrosion Cracking of Selected Copper Alloys Exposed as U-Bends at Various Atmospheric Sites Alloy 110 194 220 230 260 353 353 422 425 425 443 443 510 638 654 669 688 706 715 722 752 752 762 770 Condition Hard Hard Extra hard Cold rolled Hard Extra hard Cold rolled Extra hard Cold rolled Extra hard Cold rolled Cold rolled Extra hard Extra hard Cold rolled Cold rolled Cold rolled Cold rolled Cold rolled Cold rolled Extra hard Cold rolled Extra hard Extra hard 50% 50% 40% 10% 40% 25% 20% 40% 50% 50% 10% 50% Alloying elements (wt%) New Haven Daytona Beach East Alton 99.5% Cu 2.4% Fe, 0.13% Zn, 0.04% P 10% Zn 15% Zn 30% Zn 35% Zn, 2% Pb 35% Zn, 2% Pb 12% Zn, 1% Sn, 0.2% P 9.3% Zn, 2% Sn 2% P 9.3% Zn, 2% Sn, 0.2% P 28% Zn, 1% Sn, 0.04% As 28% Zn, 1% Sn, 0.04% As 5% Sn, 0.2% P 2.8% Al, 1.8% Si, 0.4% Co 3% Si, 1.6% Sn 15% Zn, 12% Mn 23% Zn, 3.4% Al, 0.4% Co 10% Ni, 1.4% Fe 30% Zn, 0.5% Fe 17% Ni, 0.75% Fe, 0.5% Cr 17% Zn, 18% Ni 17% Zn, 18% Ni 30% Zn, 12% Ni 27% Zn, 18% Ni NF a NF NF NF 133 856 79 NF NF NF NF 69 NF NF NF 11 2080 NF NF NF NF 3100 667 396 NF NF NF NF 140 NF 1256 NF NF NF NF 61 NF NF NF 7.5 NF NF NF NF NF NF 1105 NF NF NF NF NF NF NF NF NF NF NF NF NF NF NF 9.5 NF NF NF NF NF NF NF 146 Source: Ref. 19. a NF ⫽ no failures. dried. The staining occurs under water drops, which are the last to evaporate. The situation contrasts with tarnishing in bulk water, which only occurs slowly because the tarnishing process requires diffusion of oxygen to the metal surface. Under a water drop, the diffusion distance of the oxygen from the air to the metal surface is small. The rates of tarnishing are then much higher, and brown stains of copper oxides can result in times of only a few minutes. Under conditions of high humidity and temperatures, such rapid tarnishing can also occur even if the metal is originally dry. On cleaned dry parts held in a barrel, moisture can condense out when the temperature falls, as, for example, during the night. This water is trapped at points where components are in contact and held there by capillary action. Rapid tarnishing, because of the ease of diffu- 136 Brock sion of oxygen, through the thin water layers results. Similar rapid tarnishing can occur between two alloy strips when one is placed on top of the other. Moisture condensing from the atmosphere is trapped in the space between the two strips and drawn into the center of the strip by capillary action. During severe conditions of humidity and temperature, tarnishing can be such as to result in the formation overnight of violet films on the surfaces of the strips which are in contact, with no tarnishing on surfaces not in contact. The rapid tarnishing described occurs on all copper alloys and the rate of tarnishing is worse if there are contaminants on the alloy surface, such as those resulting from ineffective rinsing following an acid clean. The tarnishing can be reduced by maintaining a low ambient humidity and temperature in the area where parts are dried and stored. Significant decrease in the tarnishing rates can also be obtained by final rinsing in water containing 0.1–0.5 wt% benzotriazole (BTA). Such treatment results in the formation of a protective film of a BTA– ˚ thick. This serves to inhibit attack of the copper complex only some 20–40 A alloy surface. VI. SOIL CORROSION Copper generally has excellent resistance to soil-side corrosion. The degree of attack varies with the nature of the soil. Thus, Gilbert (39) examined the corrosion of copper buried in soils at seven sites in England for times of up to 10 years. He showed that the most corrosive soils were a wet acid peat with a pH of 4.3 and a moist acid clay with a pH of 4.6. The corrosion rates in these averaged up to 260 µin./year with a pitting rate of up to 1.8 mils/year. In the less corrosive soils, the general corrosion rate ranged from 2 to 10 µin./year and there was no pitting. In a further test, a 5-year exposure of phosphorus-deoxidized copper in cinders resulted in high corrosion rates of up to 260 µin./year, with pitting at a rate of 12.6 mils/year. These high corrosion rates were shown to be due to the presence of sulfides and sulfate-reducing bacteria. Logan and Romanoff (40) evaluated samples exposed for up to 14 years in 14 different soils in the United States. The greatest attack was observed in soils where the backfill contained cinders or had high organic or inorganic acid content. Losses in wall thickness ranged from 0.2 to 2 mils and pit depths were up to 51 mils. From the above and work conducted by Denison (41), Meyers and Cohen (42) described conditions which could render soils corrosive to copper. These included elevated sulfate and chloride contents, poor drainage, inorganic and organic acids, cinder fills, sulfate-reducing bacteria, and ammoniacal compounds. Other factors affecting the corrosion behavior are the aeration characteristic of the soil, differential aeration, and stray currents. Corrosion of Copper and Its Alloys 137 Fig. 8 Weight loss–time plots for selected copper-alloy tubes resulting from New Haven tap water flowing at 8 ft/s at ambient temperature. Cohen and Brock (43) recently described the soil-side corrosion of copper water service lines removed from various streets in Billings, Montana after service in the range 10–70 years. Their findings attested to the excellent resistance of copper in the alluvial soil characteristic of this location. In some instances, tubes in use for 70 years had essentially no evidence of corrosion on the soil side. The degree of corrosion varied from site to site, but in the worst case, pits only 11 mils deep were found, and these were on a tube which had been in service for 35 years. VII. WATER CORROSION A. Potable Water Copper and its alloys generally have an excellent resistance to corrosion in freshwater and are widely used in water distribution systems and as plumbing tubes. In the United States, some 375,000 miles of plumbing tube is installed each year. Mostly, this is in the form of phosphorus deoxidized, alloy C12200, tubes. In Ref. 43, these alloy tubes were shown to have suffered little if any water-side attack even after 70 years of service. The corrosion rates of copper alloys in water are also a function of alloy type. Figure 8 shows weight loss–time plots for three alloy tubes exposed to New Haven tap water flowing at 7 ft/s under ambient conditions (19). The rate of corrosion decreases with time due to the formation of protective copper oxide films. Table 8 summarizes the corrosion rates in New Haven water after 1 year. The lowest corrosion rate was observed with alloy C19400 containing 2.4% Fe 138 Brock Table 8 Corrosion Rates of Copper Alloys After 1 Year in Flowing New Haven Tap Water at 7 ft/s Alloying elements (wt%) Corrosion rate (mils/year) 0.02% P 2.4% Fe, 0.13% Zn, 0.04% P 15% Zn 20.5% Zn, 2% Al 10% Ni, 1.4% Fe 0.32 0.1 0.4 0.2 0.27 Alloy C12200 C19400 C23000 C68700 C70600 Source: Ref. 19. and the highest for alloy C230 with 15% Zn. The corrosion rates are low and all less than 0.5 mils/year. Despite the outstanding corrosion resistance, there are isolated instances where corrosion can lead to problems. Pitting of copper tubes, although infrequently found, can occur in certain aggressive waters and lead to early failure. In water at temperatures above 140°F, pitting is particularly unusual. Waters which promote such pitting are soft, have a pH less than 7.4, and have a ratio of bicarbonate to sulfate ion content of less than 1 (44). Failures have occurred in Canada, Sweden, Europe, and the United Kingdom. The author has seen rapid failure of copper tubes in hot, soft well water in Connecticut (U.S.A.). Pits which form in hot water are deep and narrow and contain cuprous oxide. They are generally capped by black or greenish-black mounds of copper oxide and basic copper sulfate. The pits are often surrounded by a deposit high in alumina. In some hot waters, the presence of manganese can promote pitting and form somewhat larger pits which are surrounded by black deposits of manganese dioxide. Cold-water pitting has been widely described in the literature, the studies mainly relating to failures in the United States, the United Kingdom, and Belgium. It requires an aggressive water. Such pitting typically arises in hard waters from deep wells. The pits are usually hemispherical in nature and form under green mounds of malachite, a basic copper carbonate. When such nodules are removed, a brown film of cuprous oxide is revealed with a central hole (Fig. 9a). Close examination of this reveals that it has an underlying crystalline nature (Fig. 9b). In cross section (Fig. 9c), this oxide is seen to form a membrane which covers the underlying pit. The base of the pits are filled with copper chloride. Lucey (45) proposed a mechanism by which such pits are formed. The copper oxide membrane is critical to his theory. He proposed that the copper oxide membrane initially promotes the formation of pockets of copper chloride adjacent to Corrosion of Copper and Its Alloys 139 the metal surface. In continuing pit growth, the membrane then separates the anodic process within the pit from the cathodic process of oxygen dissolution on the water side of the membrane. His generally accepted mechanism clearly requires the presence of chloride and bicarbonate ions in the water, as well as dissolved oxygen. Cohen and Lyman (46) analyzed 65 waters in the United States where pitting had occurred, and they showed that they typically contain over 5 ppm carbon dioxide, have a pH value in the range 7.0–7.8, contain 10–12 ppm oxygen, and have a sulfate content generally three to four times that of the chloride ion content. These findings are in accord with the work of O’Brecht et al. (47), who showed that aggressive waters contained high concentrations of dissolved carbon dioxide and oxygen, a pH in the range 7.0–7.8, and a sulfate-to-chloride ratio of from 3–4 to 1. Following analysis of many aggressive waters in the United Kingdom, Lucey (48) produced a nomogram, the use of which permits termination of the pitting propensity of the water as determined from its sulfate, sodium, chloride, and bicarbonate ion content, the concentration of dissolved oxygen, and the pH. The presence of carbon films can promote pitting in aggressive waters (49,50). Such films can be present on a half-hard tube as a result of the decomposition of drawing oil during annealing operations. The films can promote pitting by isolating pockets of copper chloride adjacent to the alloy surface and enhance pitting by providing a large cathode-to-anode ratio. In certain potable waters, corrosion of plumbing fixtures, which are usually fabricated from leaded brass alloys, can occur. The factors which promote dezincification are high temperatures, water stagnation, the presence of crevices where aeration is restricted, a high ratio of chloride ion to carbonate ion, and a relatively high pH (2). In terms of plumbing fixtures, such an attack causes problems not only because of metal degradation but also because of the formation of voluminous corrosion products which can cause blockages of valves and freezing of valve stems. B. Freshwater Cooling Systems The excellent corrosion resistance of copper alloys in water results in their widescale utilization in power utility condensers and heat exchangers where the source of water is from rivers or lakes. Alloys typically used for condenser applications include C14200, C19400, C44300, C68700, C70600, and C71500. In systems in which the cooling water is circulated through cooling towers, evaporation, blowdown, and makeup of the water results in an increase in concentration of the salts initially present in the water. The waters can then become more corrosive. Copper–nickel alloys find more application in such systems. 140 Brock Corrosion of Copper and Its Alloys 141 Fig. 9 Micrographs showing characteristics of pits formed on copper alloys in potable water: (a) perforated membrane of cuprous oxide exposed after removal of corrosion product nodule at 25⫻; (b) details of a membrane revealing its crystalline nature at 1000⫻; (c) cross section through small nodule revealing pit, corrosion products in pit, oxide membrane, and overlying nodule at 200⫻. C. Saltwater Copper alloys have a high resistance to corrosion in seawater. Although copper itself has good resistance, the copper–nickel alloys are among the most resistant. This is illustrated in the weight-loss results shown in Table 9 obtained for panels of selected copper alloys exposed for various times on racks below the low-tide level at Daytona Beach (19). These results are consistent with those obtained from much longer immersion in tidal seawater (51). In these, the corrosion rates of alloys C70600 and C71500 after immersion for 14 years were shown to be only 1.1 and 0.8 µm/year. The good corrosion resistance of copper alloys in saltwater or seawater has resulted in their wide-scale use in ships and in tidal power station condensers. Alloys used in such applications include the inhibited Admiralty alloys C44300, C44400, and C44500, and alloys C61300, C68700, C70600, C71500, and C71640. Instantaneous corrosion rates and weight losses of selected copper alloys after 1 year obtained from laboratory studies in 3.5% sodium chloride solution 142 Brock Table 9 Weight Loss of Copper Alloys After Immersion Below Low-Tide Level for Various Times at Daytona Beach 60 Days 156 Days 365 Days 99.5% Cu 2.4% Fe, 0.13% Zn, 0.04% P 15% Zn 30% Zn 10% Ni, 1.4% Fe 27% Zn, 18% Ni 10.6 9.6 8.9 17.2 3.6 9.7 14.6 11 102 15.6 3.4 14.2 19.7 16 16.1 21.9 5.1 20.9 Alloy C11000 C19400 C23000 C26000 C70600 C77000 Weight loss (mg/cm 2) Alloying elements (wt%) Source: Ref. 19. at 40°C and flowing at 5 ft/s are shown in Table 10 (19). The various weight losses reflect the more rapid rate of corrosion of the alloys which occurs during the early stages of the tests. The corrosion resistance of copper alloys in nonpolluted seawater results from the formation of protective films over the surface of the alloy. These have been shown to be cuprous oxide with outer layers enriched in iron and nickel (52). Analysis of the protective films formed on a copper–nickel condenser tube removed from tidal power stations alloys confirms these findings (19). Highresolution scanning electron microscopy of cross sections of such films always reveals an inner layer of cuprous oxide with layers high in nickel and then iron Table 10 Corrosion Rates and Weight Losses of Copper Alloys After 1 Year with 3.5% Sodium Chloride at 40°C Flowing at 7 ft/s Alloy Alloying elements (wt%) Corrosion rate (mils/year) Weight loss (mg/cm 2) C12200 C19400 C44300 C68700 C70600 C71500 0.02% P 2.4% Fe, 0.13% Zn, 0.04% P 28% Zn, 1% Sn, 0.04% As 20.5% Zn, 2% Al 10% Ni, 1.4% Fe 30% Ni, 0.5% Fe 0.6 ⬍0.1 0.45 ⬍0.1 ⬍0.1 ⬍0.1 38.6 16.0 10.3 4.0 4.7 3.8 Source: Ref. 19. Corrosion of Copper and Its Alloys 143 oxides above this. Above these, are a layer of paratacamite, Cu2(OH)Cl, and, invariably, a layer of material deposited from the water, which usually has a high silica content often with significant concentrations of oxides of iron. Examination of Alloy C687 condenser tubes from a coastal power station revealed that some 20% of the surface was covered with an orange layer containing iron, copper, oxygen, and sodium (53). The remaining surface was covered with a double layer, with the outer layer being a porous material consisting mostly of iron oxide and an inner later consisting of hydrotacite, Mg 6 Al 2 (OH) 16 CO3 4H2O, and lesser amounts of paratacamite. Further work on tubes removed from many power stations revealed that there was always an outer layer of γ-FeOOH with a thin underlying layer containing Mg, Al, Zn, and Cu (54). Selection of alloys for seawater condenser tubes is based on factors other than their intrinsic corrosion resistance. An important consideration is the resistance of the alloy to erosion–corrosion resistance of the material. This determines not only the maximum coolant velocity but also the resistance to increased turbulence round partial blockages in the tubes. Acceptable maximum velocities obtained from both laboratory and service performance are listed in Table 11 (55). The high resistance of copper–nickel alloys to erosion–corrosion by seawater was demonstrated in the early pioneering work of Stewart and LaQue (56). They showed that at velocities of 30 ft/s, the erosion–corrosion resistance of Cu–10% Ni alloys increased with increasing iron content and leveled off at the 1.4% iron Table 11 Accepted Maximum Tubular Design Velocities for Some Copper Alloys Alloying elements (wt%) Maximum velocity (ft/s) 0.02% P 28% Zn, 1% Sn, 0.04% As 5% Al 7% Al, 0.3% Sn 20.5% Zn, 2% Al 1.5% Si 10% Ni, 1.4% Fe 30% Ni, 0.5% Fe 17% Ni, 0.75% Fe, 0.5% Cr 2–3 4–6 9 9 8 3 10–12 15 30 Alloy C12200 C44300 C60800 C61300 C68700 C65100 C70600 C71500 C72200 Source: Ref. 25. 144 Brock content. Some of this resistance was lost when the alloy was annealed. This precipitates the iron out in an iron–nickel–copper phase. These findings were instrumental in setting the iron content of Alloy C70600 close to 1.4% with magnetic permeability measurements being conducted to ensure that the iron is essentially retained in solid solution. Similar increases in erosion–corrosion resistance of copper–30% nickel were observed when the iron content was increased from 0.05% to 0.46%. Increasing the iron content above these levels can result in a increased tendency for pitting corrosion (57). The presence of 2% Mn and 2% Fe in an alloy with 15% Ni, C71640, and the Cr addition in Alloy C72200 also leads to high resistance to erosion–corrosion and in preventing attack in severe conditions where there is entrained sand (57). VIII. FOULING RESISTANCE During seawater service, fouling of materials can occur. This phenomenon consists of the attachment of marine organisms on the surface of the materials. These organisms include algae, sponges, barnacles, oysters, and mussels. Copper alloys have a high resistance to such fouling, with Alloy C70600 being the best and superior in this respect to Alloy C71500 and Alloy C68700 (58). Studies of the fouling resistance of copper alloys have attributed the affect to poisoning of the organisms by the slow release of copper ions from the alloy surface (59) and to the toxicity of cuprous oxide to marine organisms as well as the sloughing off of outer corrosion product layers of basic copper oxides (60). Marine fouling can result in the buildup of substantial layers of marine organisms on boat hulls, resulting in damage and added resistance to movement through the water. The resistance of copper–nickel alloys to biofouling has resulted in their successful use for hulls of yachts and shrimp trawlers (61). In tidal power stations, copper–nickel alloys have proved useful as seawater in the condenser screens because of the resistance to both corrosion and biofouling. IX. EFFECT OF POLLUTION The high corrosion resistance of copper alloys in waters is adversely affected when pollutants are present—in particular, sulfides. The sulfides are introduced into the water by sulfate-reducing bacteria which under anaerobic conditions reduce sulfate ions to sulfides. Alternatively, sulfides can be introduced into waters by the decomposition of plant or animal matter. The decrease in corrosion resistance is associated with the incorporation of sulfides into the corrosion product Corrosion of Copper and Its Alloys 145 films. These render the films far less protective to subsequent corrosion. Pitting of copper alloys is typical in sulfide-polluted conditions. The sensitivity of the corrosion resistance to sulfide ions is remarkable, and sulfides at the level of only 0.01 ppm can promote attack on Alloy C706 (62). Other work has demonstrated that the adverse affect of sulfides increase with increasing coolant velocity (63). Polluted waters can lead to early failures of condensers. Elimination of the source of the pollutant is an obvious step in preventing this type of attack. Decaying plant and animal life can be prevented from entering condensers by the use of screens and filters. Aeration of the water is also beneficial in removing sulfides. The intentional formation of hydrated iron oxides is also beneficial in decreasing the rate of attack by polluted waters. Iron can be introduced into waters either by an iron anode or by intermittently introducing ferrous ions into the water by the addition of ferrous sulfate (64). The protective films are then deposited over the surface of the existing corrosion product films. X. STEAM In the absence of carbon dioxide, ammonia, and oxygen, copper alloys are resistant to attack by steam. This property is important in the wide use of copper alloys in condenser applications. At high steam velocities, such as may be seen where steam enters a condenser, erosion of copper alloys can occur. Additionally, if the steam is wet, then water droplets impinging on the metal can promote severe attack in the form of a high frequency of narrow, deep pits. This is illustrated in Fig. 10 for an Admiralty alloys tube which had been exposed to such an impingement attack (19). The attack can be averted by incorporation of appropriate baffles to prevent the water droplets and steam from directly impinging on the tubes. XI. AMMONIACAL SOLUTIONS Concentrated ammoniacal solutions are corrosive to copper alloys and, as described previously, can promote stress-corrosion cracking. In service conditions copper-alloy condenser tubes often are exposed to condensate which contain low levels of ammonia. The ammonia originates from the decomposition of oxygen scavengers such as hydrazine or morpholine. The concentrations of ammonia are highest in the air-removal section of the condensers. Although the concentrations of ammonia are only a few ppm, its presence together with oxygen can promote attack. Copper–nickel alloys are most resistant to attack by ammoniated conden- 146 Brock Fig. 10 Micrograph (at 100⫻) of pits on condensate side of an alloy C706 condenser tube resulting from impingement of water droplets in wet steam. sate. This is supported by the data for various copper alloys under both laboratory and field test conditions, as shown in Table 12 (65). XII. CORROSION IN OTHER ENVIRONMENTS The preceding sections describe the behavior of copper alloys in environments in which they are most used. The metals are also used in a wide range of chemical equipment, including pipelines, fractionating columns, heat exchangers, and condensers. Space does not permit detailed descriptions of these. Such information is available in a number of publications (66,67). The subsequent subsections give a general description of the behavior of copper and its alloys in chemical environments. A. Acid Solutions All copper alloys are attacked by oxidizing acids such as nitric acid and strong sulfuric acid. In deaerated nonoxidizing acids, copper is essentially immune to corrosion. The presence of dissolved air in the acid will, however, result in attack Corrosion of Copper and Its Alloys 147 Table 12 Comparison of Field and Laboratory Data for Condensate-Side Corrosion of Copper Alloys Corrosion rate (mils/year) Alloy Plant A Plant B Plant C Laboratory C71500 C72200 C70600 C44300 0.0083 0.016 0.019 0.05 0.004 0.016 0.014 0.031 0.015 0.015 0.018 0.024 0.002 0.008 0.043 0.09 Note: Corrosion rates are after 2 years exposure with laboratory data extrapolated from 1000 h in 2 ppm ammonia at pH 9.4. Source: Ref. 19. of the metal. The rate of attack in acids is therefore proportional to the acid concentration, the temperature, the degree of aeration, and the flow rate. Under fairly mild conditions, copper alloys are successfully used for handling hydrofluoric, sulfuric, phosphoric, acetic, and other organic acids (68–71). Alloys most resistant are the tin–bronzes, the aluminum bronzes, the silicon bronzes, and the cupronickels. B. Alkali Solutions Copper alloys have good resistance to alkali solutions other than ammonium hydroxide. Copper–nickel alloys are most resistant and corrosion rates of less than 0.2 mils/year are typical of Alloy C71500 in 1N to 2N sodium hydroxide solutions. For the phosphor–bronze alloys, the corrosion rates under the same conditions are some 10 mils/year. C. Neutral Solutions Copper and its alloys are suitable for handling most neutral solutions of nonoxidizing salts. They are used for handling solutions of nitrates, sulfates, and chlorides. Solutions of oxidizing salts, such as those containing chromate, ferric, or stannic ions, can promote rapid attack. Similarly, salts of metals more noble than copper will promote attack while plating out on the metal surface. D. Organic Compounds Copper alloys are resistant to a wide range of organic compounds such as amines, ester, ethers, ketones, alcohols, aldehydes, naphtha, and gasoline. In amines, the 148 Brock corrosion resistance is significantly decreased if there is contamination by water. These conditions can also promote stress-corrosion cracking of brass. Alloys C44300 and C71500 are used in equipment for refining gasoline. Copper is extensively used for kettles in the brewing of beer and for evaporators and heating coils in the manufacture of cane and beet sugar. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. MED Turner. Proc Soc Water Treat Exam 10:162, 1961. MED Turner. 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New York: McGraw-Hill, 1950. 6 Reactive and Refractory Alloys Te-Lin Yau Te-Lin Yau Consultancy, Albany, Oregon I. INTRODUCTION This chapter covers three metals each of Group IVB [titanium (Ti), zirconium (Zr), and hafnium (Hf )] and Group VB [vanadium (V), niobium (Nb), and tantalum (Ta)] in the periodic table. Because of their similarities, these metals are referred to as the following: 1. Reactive metals, as they are highly active in the electromotive force (EMF) series. Practically, the reactivity of these metals allows them to spontaneously form protective oxide films in air and makes them corrosion-resistant materials. 2. Refractory metals, as their melting points are above the range of iron, cobalt, and nickel. Although these metals are similar in many ways, they also show very significant differences. The attractive strength-to-density ratio makes titanium an important structural material in the aerospace industry. The contrast between zirconium and hafnium in thermal neutron absorption cross section makes them complementary in nuclear applications. Vanadium is a vital alloying element to make steel and titanium strong and tough. Niobium has hot and cold applications because of its strength at elevated temperatures and superconducting characteristics at low temperatures. Tantalum is widely used in electrolytic capacitors resulting from the high dielectric constant of its surface oxide film. In addition to these standout applications, materials and corrosion professionals often specify Ti, Zr, Nb, Ta, and their alloys to handle highly corrosive environments. Moreover, current trends in the chemical process industries favor the increas151 152 Yau ing usage of these metals. These metals are highly corrosion resistant over wide ranges of media. When properly used, these metals can realize high returns-oninvestments due to low maintenance and replacement costs, improved process efficiency, added values of high-quality products, and compliance with safety and environmental protection requirements. However, there are some striking differences among these metals, too. It is essential to fully understand the capabilities of these metals. Otherwise, any misuses may result in costly mistakes. Consequently, this chapter discusses environmental effects on reactive and refractory metals. II. GENERAL CHARACTERISTICS Physical and some typical mechanical properties of unalloyed reactive and refractory metals are given in Table 1. These metals have high to very high melting points and their densities range from low to high. Their coefficients of thermal expansion are typically low. Their thermal conductivities are from better than that of type 304 stainless steel to matching that of steel. Their mechanical properties vary over a wide range depending on metallurgical states and impurities levels. The presence of interstitial impurities, such as oxygen and nitrogen, are particularly pronounced in affecting mechanical properties. In fact, oxygen can be used to modulate mechanical properties. It is important to properly specify any metal to match requirements in a specific application. There are different grades of unalloyed reactive and refractory metals by varying impurity levels. Also, there are alloys developed for improved mechanical and/or corrosion properties. Readers should review the ASTM specifications shown in Table 2 before specifying them. III. ENVIRONMENTAL EFFECTS Reactive and refractory metals are often regarded as materials that possess exceptional corrosion resistance in a wide range of environments. One can expect that these metals are similar in corrosion resistance in many environments, such as seawater. However, it is not always clear which metal is the right choice for a specific application. One should consider these metals as friendly rivals that display contrasts. Very often, the selection is made based on cost and availability. This approach may work simply because these metals are overlapping in corrosion properties. However, this approach may result in costly mistakes, too. It is important to fully understand their corrosion properties. Candidates can be selected by matching the capabilities of these metals to the characteristics of the environments. Then, the optimal choice can be made by considering other factors, such as cost, requirements of product quality, strength, and fabricability. Physical and Mechanical Properties at Room Temperature Property Atomic number Melting point (°C) Density (g/cm3) Coefficient of thermal expansion (⫻10⫺6 /°C) Thermal conductivity (W/m°C) Modulus of elasticity (GPa) Tensile strength (MPa) Yield strength (MPa) Stress-relieving temperature (°C) Ti Zr Hf V Nb Ta 22 1677 4.51 8.9 17 52 240–655 170–720 480 40 1857 6.51 5.89 22 99 165–440 170–310 540 72 2227 13.3 5.9 22.3 135 440 225 650 23 1917 6.11 9.4 31 120 200–⬎500 100–⬎200 900 41 2468 8.57 7.1 523 188 170–⬎500 75–⬎200 800 73 2996 16.69 6.6 544 185 170–285 100–170 900 Reactive and Refractory Alloys Table 1 153 154 Yau Table 2 ASTM Specifications Products Ti Zr Hf Nb Ta Plate Pipe Tube Wire Ingots Castings Fittings Forgings B-265 B-337 B-338 B-348 — B-367 B-363 B-381 B-551 B-658 B-523 B-550 B-495 B-752 B-653 B-493 B-776 — — B-737 — — — — B-393 — B-394 B-392 B-391 — — — B-708 — B-521 B-365 B-364 — — — Table 3 gives a general comparison on the corrosion resistance of reactive and refractory metals. Hafnium and vanadium are omitted because they are not being used in the chemical process industries. The major purpose of this comparison is to show that there are major differences among them. It should be noted that the magnitude of the differences depends on metallurgical and chemical factors. A. Titanium Titanium is most popular among all reactive and refractory metals. In fact, Ti is the fourth most abundant structural metal in the world. The once exotic image Table 3 General Comparison of the Resistance of Reactive Metals Condition Ti Zr Nb Ta Hydrochloric acid Sulfuric acid Nitric acid Oxidizing acids w/o Cl⫺ Oxidizing acids with Cl⫺ Acids with F⫺ Caustics Hydrogen peroxide Dry chlorine Wet chlorine Ignition in oxygen Abrasives a Excellent Very good Excellent Excellent Cautious Poor Excellent Excellent Very good Cautious Very good Very good Good Good Excellent Excellent Excellent Fair Poor Fair Excellent Excellent Excellent Good Very good Excellent Excellent Excellent Excellent Poor Poor Fair Excellent Excellent Excellent Good a Poor Poor Very good Very good Excellent Poor Poor Poor Poor Excellent Good Good Ranking (excellent being the best and poor being the worst): Excellent, very good, good, fair, and poor. Cautious means that other factors such as surface condition is important. Reactive and Refractory Alloys 155 of Ti is long gone. Titanium’s popularity can be attributed to its several attractive properties, including competitive cost, good corrosion resistance, high availability, superior structural efficiency, and great fabricability. In addition to the well-known aerospace applications, Ti is widely used in industrial, marine, and commercial applications that include pulp and paper, desalination plants, dental materials, jewelry, architectural, marine, and sporting equipment, medical implants, automotive, electrochemical anodes, food, brewing, pharmaceutical, flue gas desulfurization, steam turbines, petrochemical refineries, nuclear waste storage, metal extraction equipment, and cookware. By itself, Ti is not suitable for most of these applications because of inadequate corrosion resistance, strength, and formability. However, unlike other reactive and refractory metals, Ti can be readily alloyed to modify its properties. Basically, there are three types of titanium alloy: 1. Alpha alloys are non-heat-treatable and are very weldable. They possess low to medium strength, good notch toughness, adequate ductility, and excellent cryogenic mechanical properties. The more highly alloyed alpha and near-alpha alloys offer improved high-temperature creep strength and oxidation resistance. 2. Alpha-beta alloys are heat treatable and most are weldable. They have medium to high strength. Their hot forming properties are good. They have lower high-temperature creep strength than most alpha alloys. 3. Beta and near-beta alloys are readily heat treatable and generally weldable. They are capable of achieving high strengths. They have good creep strength to intermediate temperatures. Solution-treated alloys have excellent formability. In some cases, it takes less than 0.5% of alloying elements to significantly improve Ti’s corrosion resistance. In other cases, it takes large amounts of alloying elements to substantially increase Ti’s strength. Also, technologies such as powder metallurgy and superplasticity have been developed to fabricate Ti into complicated components. It would be too extensive for this chapter to cover all the details. Only some information relating to environmental effects will be discussed here. 1. Water and Steam Titanium and its alloys are highly resistant to water, natural waters, and steam to temperatures up to at least 316°C. They may acquire a tarnished appearance on their surfaces in hot water or steam due to the formation of a protective oxide film. This is normal and causes no alarm. Natural waters often contain contaminants, such as iron and manganese 156 Yau oxides, sulfides, carbonates, and chlorides, that do not affect Ti’s corrosion resistance. Also, chlorination treatments used to control sliming and biofouling have no adverse effects on Ti. 2. Seawater and Other Salt Solutions Seawater is a complicated corrosive. It is difficult for common metals and alloys to handle. Frequently encountered corrosion problems include general corrosion, pitting, stress-corrosion cracking (SCC), microbiologically induced corrosion (MIC), and erosion. Titanium resists corrosion by seawater, regardless of chemistry variations and pollution effects, to temperatures up to 260°C. The compatibility between Ti and seawater makes Ti a vital material in marine applications. Titanium equipment has provided reliable service for decades in the chemical, oil refining, and desalination industries. Titanium does not encounter the common problems like other metals and alloys in seawater. It is practically immune to pitting, MIC, and stress-corrosion cracking. However, some highly alloyed Ti alloys are susceptible to stresscorrosion cracking. Titanium can withstand seawater impingement and flow velocities in excess of 30 m/s. Similarly, titanium resists attack by chloride solutions and other brines over the full concentration range with pH between 3 and 11. Oxidizing chloride solutions, such as ferric and cupric chlorides, chlorites, hypochlorites, chlorates, perchlorates, and chlorine dioxide, extend titanium’s resistance to lower pH levels. Nevertheless, titanium is not perfect in hot seawater and other chloride solutions. One of the major concerns is its susceptibility to crevice corrosion within tight physical crevices. It is affected by several, often interacting, factors, including temperature, solution chemistry/pH, nature of the crevice, alloy composition, metal surface condition, and metal potential. One can apply one or more measures for preventing crevice corrosion on Ti equipment and components. These measures include selecting the right alloy (e.g., palladium-containing alloys), noble metal surface treatment, pickling for smeared surface iron particles, and avoiding incompatible gaskets/sealants. Another major concern is galvanic corrosion leading to hydrogen embrittlement. Normally, titanium is the cathode when it is in contact with common metals, such as steel and aluminum. Titanium does not corrode, but the coupling metal becomes the anode and experiences accelerated corrosion. Consequently, the excessive hydrogen generated resulting from the contact may induce hydrogen embrittlement in titanium. To avoid galvanic corrosion, it is important to use two metals that are close in the galvanic series. Other preventative Reactive and Refractory Alloys 157 measures include insulating joints, coatings, and cathodically protecting the anode. 3. Inorganic Acids Titanium is regarded as a resistant metal in oxidizing acids, such as nitric and chromic acids, over a wide range of concentrations and temperatures. It is fair in mildly reducing acids, such as sulfurous acid, but is rather poor in strongly reducing acids, such as sulfuric, hydrochloric, hydrobromic, and phosphoric acids. Still, titanium is not as corrosion resistant as zirconium and tantalum to nitric acid. Titanium resists nitric acid over a wide range of concentrations at temperatures below boiling temperatures. At boiling and above, titanium’s corrosion resistance is very sensitive to nitric acid purity. Generally, titanium is not corrosion resistant in pure nitric acid. Similarly, titanium may be attacked in the vapor of nitric acid where condensates may form. The higher the metallic ion content of the acid, the better titanium will perform. In particular, the corrosion of titanium to produce small amounts of titanium ions will result in the inhibitive effect on titanium’s corrosion in nitric acid. Therefore, titanium’s corrosion may decrease rapidly in nitric acid under a closed system. Unlike stainless steels, titanium is uniquely suitable for handling recycled nitric acid. Consequently, titanium has many commercial applications. When titanium experiences corrosion problems in nitric acid, the problems cannot be solved by switching to normally more resistant titanium alloys, such as Pd-containing alloys. These types of alloys are useful in improving Ti’s resistance in reducing acids. Because nitric acid is not reducing but oxidizing, the switch will not offer any improvement. Moreover, titanium should not be considered for handling red fuming nitric acid because of the danger of pyrophoric reactions. Titanium has limited usefulness in strongly reducing acids (e.g., up to about 7% hydrochloric or sulfuric acid at room temperature). The resistance decreases rapidly with increasing temperature. It improves when the acids contain small amounts of oxidizing impurities, such as ferric ions or chlorine. For example, the addition of 2 g/L ferric chloride will reduce the corrosion rate of Ti in 3% HCl at boiling from 14 mm/year to less than 0.01 mm/year. Fortunately, it is common to have this type of impurities in industrial acids. Titanium does have numerous industrial applications involving reducing acids, such as in the mining industry. Titanium can have various corrosion problems in handling reducing acids, such as crevice corrosion and vapor-phase attack or when the concentration gets too strong. The addition of up to 0.25% Pd to Ti will significantly improve Ti’s resistance in reducing acids. Traditionally, the Pd content is controlled in the 158 Yau 0.15% range. Still, this addition greatly increases the cost. Recently, new grades of Ti alloys have been added by just having 0.05% Pd for improved corrosion resistance with a minimum increase in cost. 4. Organic Solutions Organic acids normally are mildly reducing. Titanium shows good corrosion resistance in most organic acids and other chemicals. It would be vulnerable to corrosion in strong organic acids, such as formic, oxalic, sulfamic, and trichloroacetic acids. Generally, the presence of oxygen due to aeration and water presence improves Ti’s resistance in organic media. On the other hand, certain anhydrous organic media may attack titanium. Without oxygen, it would be difficult for Ti to maintain its passivity. For example, dry methyl alcohol can cause SCC in titanium, probably due to the breakdown of passive film by halides. Once the passive film is broken down, there will not be any repair if there is not any oxygen or water. Indeed, the combination of the absence of water and the presence of halogens or halides is the major reason that titanium experiences corrosion problems in organic solutions. The addition of 1.5% water can suppress titanium’s susceptibility to SCC in methyl alcohol. Another major problem for Ti in organic media is its susceptible to hydrogen embrittlement. Under reducing conditions, Ti will slowly absorb hydrogen even when the corrosion rate is very low. 5. Alkalis Titanium resists most alkalis except hot, strongly alkaline solutions. The major problem is the excessive hydrogen uptake and eventual embrittlement of titanium at temperatures above 80°C when the pH is at or above 12. The presence of a strong oxidizer, such as chlorine, makes Ti highly suitable for processing alkalis. In fact, Ti is a useful structural material in the dual production of caustic soda and chlorine by an electrochemical process. 6. Gases Titanium has excellent resistance to air and oxygen at temperatures up to 370°C. Above this temperature but below 450°C, titanium may form colored surface films that thicken slowly with time. Above 650°C, titanium will become brittle due to poor oxidation resistance. Scales form rapidly at 930°C. Because the oxidation is an exothermal reaction, titanium may ignite in pressurized oxygen under a confined condition. Nitrogen reacts much more slowly with titanium than oxygen. It reacts with nitrogen to form a gold-colored film starting at 540°C. Above 800°C, excessive Reactive and Refractory Alloys 159 diffusion of the nitride will occur and may cause metal embrittlement. Properly formed nitride layer can enhance Ti’s resistance to abrasives. The surface oxide film can protect Ti from absorbing hydrogen at temperatures below 80°C. Absorption of several hundred parts per million (ppm) of hydrogen results in embrittlement and the possibility of cracking under conditions of stress. The presence of moisture in hydrogen (e.g., as low as 2%) can effectively reduce hydrogen uptake. Titanium is resistant to corrosion by sulfur dioxide and water-saturated sulfur dioxide. Water-saturated sulfur dioxide may form the highly corrosive sulfurous acid that does not affect titanium. This capability allows Ti to be used in various sulfur dioxide scrubber systems. Titanium resists attack by wet and dry hydrogen sulfide. It is well known that hydrogen can induce hydrogen embrittlement on many metals and alloys. This possibility exits in Ti, too. Unlike most metals and alloys, Ti does not become brittle in wet hydrogen sulfide. However, in galvanic couples with certain metals, such as steel, the presence of hydrogen sulfide in an aqueous solution will promote hydriding in titanium. Titanium is among the most resistant metal in wet chlorine and other chlorine chemicals because of their strongly oxidizing natures. This has been the primary factor for using Ti in industrial service. Titanium equipment has been relatively free of corrosion problems for decades. However, titanium is incompatible with dry chlorine that can cause a rapid attack of Ti and may even cause ignition. As little as 1% water is sufficient for repassivation after mechanical damage to Ti in chlorine gas under static conditions at room temperature. 7. Liquid Metals Titanium has good resistance to many liquid metals at moderate temperatures. It has been used in processing molten aluminum. Rapidly flowing molten aluminum, however, will erode titanium. Also, some liquid metals, such as mercury and cadmium, can cause SCC in titanium. B. Zirconium Zirconium has the reputation of being one of the most corrosion-resistant metals. It has a very strong affinity for oxygen. It is one of very few metals that even can react with oxygen in water under highly reducing conditions to form an adherent, protective oxide film on its surface. This film is self-healing and protects the base metal from chemical and mechanical attack at temperatures to 350°C. Many engineering metals, such as iron, nickel, chromium, and titanium, produce metal ions of a variable valency. Uniquely, zirconium is predominantly quadrivalent in its oxides and many other compounds. It forms very few unstable 160 Yau compounds in which its valency is other than 4. The chemistry of zirconium is characterized by the difficulty of achieving an oxidation state less than 4. This character, along with high oxygen affinity, allows zirconium to form protective oxide films even in highly reducing media, such as hydrochloric acid and dilute sulfuric acid. Under these conditions, common metals and alloys may form subordinate oxides or other compounds of low or no protective capability. Moreover, metal ions of a constant valency imply their stability. This is an important requirement in many applications. For example, zirconium can maintain the stability of certain chemicals, such as hydrogen peroxide. Ions of a variable valency are the common decomposition catalysts for hydrogen peroxide. Another advantage is that zirconium ions are colorless. This is important when the color stability of products is a major concern. Most transition metals produce ions of different colors depending on their valency state. Protective oxide films are difficult to form on zirconium’s surface in a few media, such as hydrofluoric acid, concentrated sulfuric acid, and oxidizing chloride solutions. Consequently, zirconium is not suitable or needs protection measures for handling these media. 1. Water and Steam Zirconium has excellent corrosion and oxidation resistance in water and steam at temperatures exceeding 300°C. Zirconium has a great capability to take oxygen from water for the formation of a protective oxide film. This capability is not weakened even when zirconium is in a highly reducing medium. Most passive metals form protective oxide films in aqueous solutions only when the solutions are somewhat oxidizing. Consequently, zirconium is uniquely suitable for nuclear applications because water-cooled reactors operate with an oxygen- or hydrogencharged coolant at temperatures from 280°C to 300°C. However, corrosion and oxidation of unalloyed zirconium in high-temperature water and steam were found to be irregular. This behavior is related to variations in the impurity content in the metal. Nitrogen and carbon impurities are particularly harmful. The oxidation rate of unalloyed zirconium increases markedly when nitrogen and carbon concentrations exceed 40 and 300 ppm, respectively. The irregular behavior of unalloyed zirconium stimulated alloy development programs. Zircaloys (Zr–1.5% Sn-based alloys), Zr–2.5 Nb and Zr–1 Nb are the most important ones developed for nuclear applications because they are more reliable and predicable for use in hot water and steam in addition to being stronger. As compared to unalloyed zirconium, zircaloy-2 has an improved character in oxide formation at elevated temperatures. A tightly adherent oxide film forms on this alloy at a rate that is at first quasicubic but undergoes a transition to linear Reactive and Refractory Alloys 161 behavior after an initial period. Unlike the white, porous oxide films on unalloyed zirconium, the oxide film on zircaloy-2 remains dark and adherent throughout transition and in the posttransition region. Zircaloy-4 differs in composition from zircaloy-2 in having a slightly higher iron content but no nickel. Both variations are intended for reducing hydrogen pickup with little effect on corrosion resistance in reactor operation. For example, in water at 360°C, hydrogen pickup for zircaloy-4 is about 25% of theoretical, or less than half that for zircaloy-2. In addition, hydrogen pickup for zircaloy-4 is less sensitive to hydrogen overpressure than that for zircaloy-2. For both alloys, hydrogen pickup is greatly reduced when dissolved oxygen is present in the medium. Zr–2.5 Nb is considered to be somewhat less resistant to corrosion than the zircaloys, with exception. Nevertheless, Zr–2.5 Nb is suitable for many applications, such as pressure tubes in the primary loops of some reactors. Furthermore, the corrosion resistance of Zr–2.5 Nb can be substantially improved by heat treatments. Also, the Zr–2.5% Nb alloy is superior to zircaloys in steam at temperatures above 400°C. 2. Saltwater Zirconium has excellent corrosion resistance to seawater, brackish water, and polluted water. Zirconium’s advantages include its insensitivity to variation in factors like chloride concentration, pH, temperature, velocity, crevice, and sulfurcontaining organism. Some of the results are summarized as follows. Zirconium specimens with or without a crevice attachment were placed in the Pacific Ocean at Newport, Oregon, for up to 129 days. All welded and nonwelded specimens exhibited nil corrosion rates. Marine biofouling was observed; however, no attack was found beneath the marine organisms or within the crevices. Laboratory tests were performed on Zr 702 (unalloyed Zr) and Zr 704 (nonnuclear grade of zircaloys) in boiling seawater for 275 days and in 200°C seawater for 29 days. Both alloys were resistant to general corrosion, pitting, and crevice corrosion. Tests of U-bend specimens, with or without steel coupling, of Zr 702, nickel-containing Zr 704, and nickel-free Zr 704 were conducted in boiling seawater for 365 days. No cracking was observed during the test period. Overstressing of the tested U-bends indicated that all specimens remained ductile except for the welded nickel-containing Zr 704 with steel coupling. Steel-coupled nickel-containing Zr 704 showed much higher hydrogen and oxygen absorption and formed hydrides. Chemical analyses and metallographic examinations on other U-bends did not show evidence of hydrogen absorption and hydride formation. 162 Yau 3. Halogen Acids Zirconium resists attack by all halogen acids, except hydrofluoric acid (HF). HF vigorously attacks it at all concentrations. Acidic fluoride solutions are highly corrosive to zirconium, too. Zirconium’s corrosion resistance is not as poor in fluoride salt solutions as in HF until fluorides become HF in pH ⱕ 3 solutions. This fact is taken advantage of in preparing zirconium surfaces using mixtures of hydrofluoric and nitric acids for various fabrication steps and for improved corrosion resistance in certain nuclear and chemical applications. In recent years, it appears that the chance to have fluorides in the process media has increased somewhat. One of the possibilities is the increased usage of recycled chemicals. For example, recycled sulfuric acid may contain more than 100 ppm fluorides. When zirconium equipment faces fluoride-containing acids, inhibitors that form strong fluoride complexes should be added for protecting zirconium equipment. Effective inhibitors include zirconium nitrate, zirconium sulfate, and phosphorus pentoxide. The other halogen acids [i.e., hydrochloric (HCl), hydrobromic (HBr), and hydriodic (HI) acids] do not attack zirconium. Yet, one of the most impressive corrosion properties for zirconium is its excellent resistance in HCl at all concentrations and temperatures even above boiling. Because of its strong reducing power, it is very difficult for most metals to form protective oxide films in HCl. The presence of even a small amount of HCl in a medium may cause common metals and alloys to suffer general corrosion, pitting, and SCC. Zirconium is suitable for handling HCl at all concentrations. Moreover, zirconium is not as susceptible to hydrogen embrittlement in HCl as tantalum is. For example, tantalum became brittle when tested in 11M HCl and 11M HCl ⫹ 7% GaCl3 for 1000 h at 70°C. Under the same conditions, zirconium remained unattacked and retained 100% of its ductility. Zirconium is susceptible to localized corrosion, such as pitting, intergranular corrosion, and SCC when it is anodically polarized. Zirconium is susceptible to pitting in ⬍20% HCl, but to intergranular corrosion in ⬎20% HCl. The same types of corrosion problem may be developed in HCl when highly oxidizing ions, such as ferric and cupric ions, are present. The presence of ferric ions may polarize the zirconium surface to a potential exceeding the breakdown potential. Thus, a local breakdown of the passive surface at preferred sites occurs, and the condition favors the occurrence of localized corrosion. To eliminate preferred sites by pickling zirconium in a mixture of hydrofluoric and nitric acids can suppress the breakdown process of passive films. Alternatively, maintaining zirconium at a potential in its passive region, which is arbitrarily set at 50–100 mV below the corrosion potential, can counteract the detrimental effects resulting from the presence of ferric ions. Reactive and Refractory Alloys 163 4. Nitric Acid Nitric acid (HNO3), because of its oxidizing power, is not considered to be a difficult acid for passive metals to handle. Nevertheless, HNO3 becomes highly corrosive when its temperature is high, its concentration is too high or not high enough, or its purity is poor. The oxidizing power favors the formation of oxide films, but it may also cause the passive films to break down. Zirconium is considerably more suitable than most passive metals for handling HNO3, particularly when the acid is hot, impure, and/or variable in concentration. In certain conditions, zirconium is even more resistant than the noble metals to the acid. Zirconium’s temperature limit is somewhat higher than that of noble metals. Traces of chloride may lead to rapid attack of noble metals, but not of zirconium. The excellent corrosion resistance of zirconium in HNO3 has been recognized for more than 30 years. Zirconium resists nitric acid up to the boiling point and at 98% HNO3, and up to 250°C and at 70% HNO3. Moreover, the corrosion rates are nil in boiling 30–70% HNO3 with up to 1% FeCl3, 1% NaCl, 1% seawater, 1% iron, or 1.45% type 304 stainless steel at 205°C. These results indicate that the presence of heavy metal ions and chlorides has little effect on the corrosion resistance of zirconium. Zirconium is normally susceptible to pitting in oxidizing chloride solutions. However, the NO3⫺ ion is an effective inhibitor for the pitting of zirconium. The minimum [NO3⫺]/[Cl⫺] molar ratio required to inhibit pitting of zirconium was determined to be 1–5. Still, the presence of an appreciable amount of HCl should be avoided because zirconium is not resistant to aqua regia. The slow strain-rate technique can reveal zirconium’s SCC susceptibility in HNO3. The primary concern for using zirconium in HNO3 service would be SCC in concentrated acid. Zirconium seems to be resistant to SCC in concentrated acids when it is stressed below the yield point. Other concerns include the accumulation of chlorine gas in the vapor phase and the presence of fluoride ions. Chlorine gas can be generated by the oxidation of chlorides in HNO3. Areas that can trap gases should be avoided when Cl⫺ is present in HNO3; or, zirconium equipment should be pickled for much improved resistance to pitting in wet-chlorine-containing vapors. As indicated previously, the corrosion of zirconium in fluoride-containing acids can be controlled by adding an inhibitor, such as zirconium compounds, to convert fluoride ions into noncorrosive complex ions. 5. Sulfuric Acid Sulfuric acid (H2SO4) is the most important acid for use in the manufacture of many chemicals. For example, it is used as a dehydrating agent, an oxidizing 164 Yau agent, an absorbent, a catalyst, a reagent in chemical syntheses, and so on. The consumption of sulfuric acid indicates a nation’s industrial activity. These highly versatile capabilities can be attributed to the complicated nature of this acid. Dilute solutions are reducing, which make passive metals vulnerable to corrosion. In fact, hot, dilute solutions can be used to pickle steel and stainless steel. Solutions become increasingly oxidizing at or above 65%. The usefulness of common metals depends strongly on acid concentration, temperature, and the presence of other chemicals. Zirconium resists attack by H2SO4 at all concentrations up to 70% and at temperatures to boiling and above. In 70–80% H2SO4, the corrosion resistance of zirconium depends strongly on temperature. In higher concentrations, the corrosion rate of zirconium increases rapidly with concentration. In the range in which zirconium shows corrosion resistance in H2SO4, a passive film is formed on zirconium that is predominantly cubic zirconium oxide (ZrO2) with only traces of the monoclinic phase. Zirconium corrodes in highly concentrated H2SO4 (e.g., 80%) because loose films are formed that prove to be zirconium disulfate tetrahydrate [Zr(SO4)2 ⋅ 4H2O]. Also, at the higher acid concentrations, films that flake off are formed and are probably partly zirconium hydrides. Zirconium cannot tolerate much of strong oxidizing agents in ⬎40% H2SO4. However, in ⬍40% H2SO4, zirconium can tolerate large amounts of strong oxidizing agents. Consequently, zirconium equipment is often used in steel pickling. The presence of chlorides in H2SO4 has little effect on the corrosion resistance of zirconium unless oxidizing agents are also present. Therefore, in the presence of oxidizing agents, chloride ions should be controlled within a limit to avoid detrimental attack. Zirconium weld metal may corrode preferentially when H2SO4 concentration is approximately 55% and higher. Heat treatment at 775 ⫾ 15°C for 1 h per 25.4 mm of thickness can be applied to restore the corrosion resistance of the weld metal to that of the parent metal. However, this high-temperature heat treatment is not suitable for equipment made of zirconium/steel clad materials because of the large difference in thermal expansion coefficients between these two alloys. Heat treatment at a much lower temperature, such as 425–530°C, should be applied when there is a concern for SCC. Zirconium is susceptible to SCC in a narrow range of H2SO4, (i.e., 64–69%). For zirconium equipment, it is very important to maintain acid concentration within the applicable limit. When the limit is exceeded, zirconium may corrode rapidly. In ⬍65% H2SO4, the vapor phase is almost entirely water. However, the concentration change is negligible when a system is under a pressurized condition. Acid concentration may change significantly when, for example, the sys- Reactive and Refractory Alloys 165 tem is imperfectly sealed. In a leaking system, the acid concentration can exceed the concentration limit. Acid concentration can easily increase when the system is under vacuum operation because water vapor is continuously taken away. When the corrosion resistance limits of zirconium in H2SO4 are exceeded, a pyrophoric surface layer may be formed on zirconium under certain conditions. The pyrophoric surface layer on zirconium formed in 77.5% H2SO4 ⫹ 200 ppm Fe3⫹ at 80°C consists of γ-hydride, ZrO2, Zr(SO4)2, and fine metallic particles. The combination of hydride and fine metallic particles is suggested to be responsible for the pyrophoricity. Treating in hot air or steam can eliminate this tendency. 6. Phosphoric Acid Phosphoric acid (H3PO4) is less corrosive than other mineral acids. Many stainless alloys demonstrate useful resistance in the acid at low temperatures. As often occurs, corrosion rates of common alloys in the acid increase with increasing temperature, concentration, and impurities. Areas like the liquid-level line or the condensing zones are particularly vulnerable to attack. Zirconium resists attack in H3PO4 at concentrations up to 55% and temperatures exceeding the boiling point. Above 55% H3PO4, the corrosion rate could increase substantially with increasing temperature. The most useful area for zirconium would be dilute acid at elevated temperatures. Zirconium outperforms common stainless alloys in this area. If H3PO4 contains more than a trace of fluoride ions, attack on zirconium may occur. Because fluoride compounds are often present in the ores for making H3PO4, the use of zirconium has always been questioned. However, because P2O5 is an effective fluoride inhibitor and is usually present in large amounts in H3PO4 processes, tests should be conducted to determine zirconium’s suitability in the actual stream. Furthermore, zirconium compounds can be used to complex fluorides. 7. Other Acids Zirconium has excellent corrosion resistance in up to 30% chromic acid at temperatures to 100°C. It is not suitable for handling chrome-plating solutions that contain fluoride catalysts. Zirconium is also resistant to some mixed-acid systems. It can be in acid mixtures of sulfuric–nitric, sulfuric–hydrochloric, and phosphoric–nitric. The sulfuric acid concentration must be below 70%. Zirconium is aggressively attacked in 1:3 volume mixtures of nitric and hydrochloric acids (aqua regia). In the 1:1 volume mixture, zirconium is attacked but much slower than in the 1:3 mixture. In mixtures greater than the 3:1 ratio, zirconium is resistant. 166 Yau 8. Alkalis Zirconium resists attack in most alkalis, which include sodium hydroxide, potassium hydroxide, calcium hydroxide, and ammonium hydroxide. This makes zirconium distinctly different from other highly corrosion-resistant materials, such as titanium, tantalum, graphite, glass, and polytetrafluoroethylene (PTFE). Zirconium U-bend specimens have been tested in boiling 50% NaOH. During the test period, the concentration changed from 50% to about 85% and the temperature increased from 150°C to 300°C. The PTFE washers and tubes used to make the U-bends dissolved. However, the zirconium U-bends remained ductile and did not show any cracks after 20 days. It should be noted that stainless steel is susceptible to SCC in alkaline solutions including NaOH solutions. Zirconium coupons were tested in a white-liquor, paper-pulping solution, which contained NaOH and sodium sulfide, at 120°C, 175°C, and 225°C. All coupons showed nil corrosion rates. In the same solution, graphite and glass both corroded badly at 100°C. Zirconium also exhibits excellent resistance to SCC in simulated white liquors. 9. Salt Solutions Zirconium is resistant to most salt solutions, which include halide, nitrate, carbonate, and sulfate salts. Corrosion rates are usually very low at temperatures at least up to the boiling point. Solutions of strong oxidizing chloride salts, such as FeCl3 and CuCl2, are examples of the few exceptions. In strong oxidizing chloride solutions, zirconium’s performance is very dependent on surface condition. Zirconium becomes quite resistant to pitting when it has a good surface finish, like the pickled surface. Although zirconium has good corrosion resistance in sodium fluoride and potassium at low temperatures, resistance decreases rapidly with increasing temperature or decreasing pH. Consequently, zirconium is not ideal for handling most fluoride-containing solutions unless fluoride ions are complexed. Zirconium is considerably more resistant to chloride SCC than stainless steels are. No failure was observed in U-bend tests conducted in boiling 42% magnesium chloride (MgCl2). Another attractive property of zirconium is its high crevice corrosion resistance. Zirconium is not subject to crevice corrosion even in acidic chloride at elevated temperatures. No attack was observed on zirconium in a salt spray environment. Unlike many common metals, zirconium has very little affinity for sulfur. Zirconium–sulfur compounds form only at temperatures above 500°C. Furthermore, there is no instance of zirconium–sulfur bonds forming in aqueous systems. Hence, hydrogen sulfide (H2S), which is highly corrosive to common metals and alloys, is not expected to participate in the corrosion reactions of zirconium in sulfide-containing solutions. Zirconium coupons and U-bends were tested in nu- Reactive and Refractory Alloys 167 merous NaCl–H2S solutions at temperatures to 232°C. No general corrosion, pitting, crevice corrosion, or SCC was observed. 10. Organic Solutions Zirconium has excellent corrosion resistance in most organic solutions, except certain chlorinated compounds. It has been extensively tested in organic-cooled reactors where the coolant consisted of mixtures of high-boiling aromatic hydrocarbons (e.g., terphenyls). These coolants are noncorrosive to zirconium. However, early experiments in the organic coolants indicated that hydriding was a major concern. It was found that chlorine impurity in the organic coolants was the major cause of gross hydriding. Elimination of the chlorine and maintenance of a good surface oxide film by ensuring the presence of adequate water (⬎50 ppm) alleviates the hydriding problem. Indeed, the combination of the absence of water and the presence of halogens or halides is the major reason why zirconium experiences corrosion problems in organic solutions. For example, the addition of some water can suppress zirconium’s susceptibility to SCC in alcohol solutions with halide impurities. On the other hand, zirconium shows excellent corrosion resistance in certain chlorinated carbon compounds (e.g., carbon tetrachloride and dichlorobenzene) at temperatures up to 200°C. From a corrosion point of view, organic halides can be classified into three groups: water soluble, water insoluble, and water incompatible. Water-soluble halides, such as aniline hydrochloride, chloroacetic acid, and tetrachloroethane, are not corrosive to zirconium. They may become more corrosive when the water content is low and/or zirconium is highly stressed. More active halides, such as dichloroacetic and trichloroacetic acids, are more corrosive to zirconium. It is suspected that these halides may attack zirconium or intermetallic compounds at grain boundaries to form organometallic compounds. It should be noted that certain organic compounds, such as alkyl and aryl halides, are the common ones that react with most metals, including noble metals, to form organometallic compounds. These reactions can be suppressed when there is some water present in the media. Consequently, water addition and/or stress relieving would be effective in preventing the corrosion of zirconium in watersoluble halides. However, water addition may increase the corrosivity of many organic solutions toward common metals and alloys, but it seems to be always beneficial to zirconium. Water-insoluble halides, such as trichloroethylene and dichlorobenzene, are not corrosive to zirconium, probably because of their stability. They will not dissolve in water and they will not exclude water, either. They and water can be physically mixed together. Water-incompatible halides, such as acetyl chloride, may be highly corro- 168 Yau sive to zirconium. They are not stable. They react violently with water. There is no chance for water to be present in this type of halides, which are the most undesirable organics for zirconium, and possibly other metals, to handle. 11. Gases Zirconium forms a visible oxide film in air at about 200°C. The oxidation rate becomes high enough to produce a loose, white scale on zirconium at temperatures above 540°C. At temperatures above 700°C, zirconium can absorb oxygen and become embrittled after prolonged exposure. Zirconium reacts more slowly with nitrogen than with oxygen because it has a higher affinity for oxygen than for nitrogen and it is normally protected by a layer of oxide film. Once nitrogen penetrates through the oxide layer, it diffuses into the metal faster than oxygen because of its smaller size. Clean zirconium starts the nitride reaction in ultrapure nitrogen at about 900°C. Temperatures of 1300°C are needed to fully nitride the metal. The nitriding rate can be enhanced by the presence of oxygen in the nitrogen or on the metal surface. The oxide film on zirconium provides an effective barrier to hydrogen absorption up to 760°C, provided that small amounts of oxygen are also present in hydrogen for healing damaged spots in the oxide film. In an all-hydrogen atmosphere, hydrogen absorption will begin at 310°C. Zirconium will ultimately become embrittled by forming zirconium hydrides when the limit of hydrogen solubility is exceeded. Hydrogen can be effectively removed from zirconium by prolonged vacuum annealing at temperatures above 760°C. The corrosion resistance of zirconium and its alloys in steam is of special interest to nuclear power applications. They can be exposed for prolonged period without pronounced attack at temperatures up to 425°C. In the 360°C steam, up to 350 ppm chloride and iodide ions, 100 ppm fluoride ions, and 10,000 ppm sulfate ions are acceptable for zirconium in general applications but not in nuclear power applications. Zirconium is stable in NH3 up to about 1000°C, in most gases (CO, CO2, and SO2) up to about 300–400°C and in dry halogens up to about 200°C. At elevated temperatures, zirconium forms volatile halides. Depending on the surface condition, zirconium may or may not be resistant in wet chlorine. Zirconium is susceptible to pitting in wet chlorine unless it has a properly cleaned surface. 12. Molten Salts and Metals Zirconium resists attack in some molten salts. It is very resistant to corrosion by molten sodium hydroxide to temperatures above 1000°C. It is also fairly resistant to potassium hydroxide. The oxidation properties of zirconium in nitrate salts are similar to those in air. Reactive and Refractory Alloys 169 Zirconium resists some types of molten metals, but the corrosion resistance is affected by trace impurities, such as oxygen, hydrogen, or nitrogen. Zirconium has a corrosion rate of less than 25 µm/year in liquid lead to 600°C, lithium to 800°C, mercury to 100°C, and sodium to 600°C. The molten metals known to attack zirconium are zinc, bismuth, and magnesium. C. Hafnium Hafnium and zirconium are chemically very similar. They occur naturally in ores. However, hafnium has a high thermal neutron absorption cross section in contrast to the very low one of zirconium. Because of this dramatic difference, they have to be separated from each other for nuclear applications but not corrosionresistant applications. Similar to zirconium, hafnium has excellent corrosion resistance to many media, including hydrochloric, nitric, and nonconcentrated sulfuric acids and alkalis. In fact, hafnium is superior to zirconium in corrosion resistance in water, steam, molten alkali metals, and air. Hafnium is soluble is hydrofluoric and concentrated sulfuric acids and aqua regia. It begins to react slowly with air or oxygen at about 400°C and nitrogen at about 900°C and rapidly with hydrogen at about 700°C. An interesting capability for hafnium is the combination of high neutron absorption and excellent corrosion resistance to nitric acid. This capability makes hafnium a uniquely reliable material for use in spent nuclear fuel reprocessing plants. Hafnium is being used not just for its corrosion resistance but also for preventing criticality. D. Vanadium Compared to other reactive and refractory metals, vanadium is inferior in aqueous corrosion resistance. There is no known application for vanadium in the chemical process industries. There is no established ASTM specification for vanadium, either. Vanadium resists attack by oxygen, nitrogen, and hydrogen at ambient temperatures. It oxidizes in air at different temperatures to form various oxides (i.e., trioxide, tetroxide, and pentoxide). It reacts with chlorine at temperatures greater than 180°C. Vanadium is resistant to seawater, reducing acids, such as hydrochloric and dilute sulfuric acids, and to alkaline solutions. It is poor in oxidizing acids, such as nitric and concentrated sulfuric acids, and hydrofluoric acid. Vanadium has one important corrosion property: its resistance in liquid metals including lithium and sodium. This, combined with its neutron economy 170 Yau and high-temperature strength, makes vanadium a leading candidate for first wall and blanket structures of liquid–metal cooled-fusion energy systems. E. Niobium The corrosion properties of niobium are often compared to those of tantalum. Like tantalum, niobium forms very stable oxide films under anodic conditions. The films are high in dielectric constant and breakdown potential. These properties, coupled with its excellent electrical conductivity, have led niobium to be used as a substrate for platinum-group metals in impressed-current cathodic protection anodes. Consequently, it would be great if one can cut costs and weights by replacing tantalum with niobium in anticorrosion applications. However, in most cases, niobium has its own unique capabilities and is not interchangeable with tantalum in severe-corrosion-resistant applications. Niobium resists most inorganic and organic acids, except hydrofluoric acid, at all concentrations and temperatures below 100°C. It is especially resistant under oxidizing conditions, such as strong sulfuric acid containing ferric or cupric ions. It has excellent corrosion resistance in most salt solutions, except those that hydrolyze to form alkalis. At room temperature, niobium is resistant to sulfuric acid at all concentrations up to 95%. The corrosion resistance decreases rapidly with increasing temperature. Compared to tantalum, niobium is much less useful in sulfuric acid applications. Niobium is excellent in oxidizing acids like nitric acid. However, niobium is susceptible to hydrogen embrittlement in reducing acids, such as hydrochloric acid. Hydrogen embrittlement can be prevented by converting the reducing condition into an oxidizing condition. For example, niobium becomes corrosion resistance in mixtures of nitric and hydrochloric acids. Like other reactive and refractory metals, niobium is not corrosion resistant in hydrofluoric acid. However, niobium has an adequate resistance in certain fluoride-containing contaminated acids. For example, niobium is suitable for handling chromium plating solutions when small amounts of fluorides are added as a catalyst. Similarly, niobium can be used in the reprocessing of spent chromium plating solutions. The presence of a fluorides makes tantalum unsuitable. At ambient temperatures, niobium is resistant in alkaline solutions. However, at higher temperatures, niobium’s corrosion rates may stay low in alkaline solutions, but it may become brittle even in dilute solutions of sodium hydroxide or potassium hydroxide. Niobium oxidizes easily in air. The oxidation becomes noticeable as the temperature approaches 200°C. It becomes rapid when the temperature exceeds 500°C. In pure oxygen, the attack is catastrophic at 390°C. Oxygen diffuses freely through the metal and this causes embrittlement. Still, within the temperature limit, niobium is much more resistant to ignition than titanium in pure oxygen. Reactive and Refractory Alloys 171 To reduce the cost, a Ti–Nb alloy can be used in ignition-sensitive areas. For example, the Ti–45% Nb alloy has found applications that employ dilute sulfuric acid and pressurized oxygen, as in the gold mining industry. Below 100°C, niobium is inert in most common gases, including bromine, chlorine, hydrogen, nitrogen, and sulfur dioxide (wet or dry). The resistance to wet and dry chlorine is particularly interesting in corrosion-resistant applications. Niobium reacts with nitrogen above 350°C, with steam above 300°C, with chlorine above 200°C, and with carbon dioxide, carbon monoxide, and hydrogen above 250°C. In the absence of nonmetallic impurities, such as gases, niobium has excellent corrosion resistance in liquid metals. The temperature limits for niobium in liquid metals are 510°C for bismuth, 400°C for gallium, 850°C for lead, 1000°C for lithium, sodium, potassium, and sodium–potassium alloys, 600°C for mercury, 850°C for thorium–magnesium eutectic, 1400°C for uranium, and 450°C for zinc. The excellent resistance of niobium in sodium vapor leads to the use of the Nb–1 Zr alloy as the end caps in high-pressure sodium vapor lamps used for highway lighting. Furthermore, liquid metals are excellent heat-transfer media; they are ideal for use in compact thermal systems, such as fast breeder reactors, reactors for space vehicles, and fusion reactors. F. Tantalum Tantalum has the reputation to be one of the most versatile corrosion-resistant materials. Like glass, tantalum is inert in most inorganic and organic solutions at temperatures to at least 150°C. The exceptions to this include fluorides, oleum, oxalic acid, and strong alkalis. In fact, tantalum parts are often used to repair glass-lined equipment because of their close corrosion properties and thermal expansion coefficients. Of course, unlike glass, tantalum can withstand thermal shocks and has a thermal conductivity comparable to carbon steel. As shown in Table 1, tantalum is quite low in strength at room temperature. The strength decreases quickly with increasing temperature. For the strength requirement at elevated temperatures, a tantalum alloy, such as Ta–2.5% W or Ta– 10% W, can be used. These alloys are as corrosion resistant as tantalum in all environments. One of the most important corrosion properties for tantalum is its resistance to the complicated sulfuric acid. Sulfuric acid changes from the highly reducing character of dilute solutions to the strong oxidizing character of concentrated solutions. Even many high-performance alloys only have very restrictive usefulness in sulfuric acid. It is unusual that tantalum resists sulfuric acid at all concentrations up to 98% and temperatures to at least 200°C. Yet, the Ta–2.5% W alloy has been shown to be even more corrosion resistant than tantalum in concentrated sulfuric acid. Due to the presence of sulfur trioxide, tantalum cor- 172 Yau rodes in fuming sulfuric acid even at room temperature. Tantalum equipment would be reliable for use in processes such as the recovery of strong sulfuric acid, the nitration, and thermal cracking of organics. Tantalum is practically inert to nitric acid, including the fuming grade, and temperatures up to 300°C. Its excellent corrosion resistance, high melting point, and high thermal conductivity make tantalum one of the most suitable materials for handling fuming nitric acid. Fuming nitric acid is not just highly corrosive but also strongly oxidizing for inducing ignition on metals. Tantalum resists hydrochloric acid at all concentrations up to 190°C, although above 25%, the corrosion rate starts to rise noticeably. In addition, the entry of hydrogen in concentrated acid may cause embrittlement in tantalum. However, the presence of small amounts of oxidizing impurities will greatly reduce hydrogen absorption to avoid hydrogen embrittlement. Also, hydrogen embrittlement can be avoided by embedding platinum particles onto tantalum’s surface in area ratios of 1: 1000 platinum to tantalum. The embedding method can be plating or welding. The low hydrogen overvoltage and cathodic character of platinum makes it very easy for atomic hydrogen to form molecular hydrogen. Thus, absorption of hydrogen into the tantalum is prevented. Tantalum is useful in handling phosphoric acid at concentrations up to 85% and temperatures up to 200°C, provided that the fluoride impurity is very low. Commercial grades of phosphoric acid may contain more than 100 ppm fluorides. When tantalum equipment is involved, the fluoride content should not exceed 5 ppm. As discussed previously, some metals, including titanium and zirconium, are vulnerable in organic halides in the absence of water and oxygen. Tantalum seems to be much less vulnerable than most metals under the same conditions. Without the interference of water and oxygen, organic halides may attack metals by forming organometallic compounds. Maybe, because of its large atomic size, it would be more difficult for organic halides to incorporate tantalum into forming a new compound. Tantalum is fairly resistant to dilute alkaline solutions. It is attacked by concentrated alkaline solutions, even at room temperature. One may not consider tantalum for caustic services. However, caution should be exercised when strong alkaline solutions are used to clean tantalum equipment. Tantalum is excellent for handling wet or dry halogens at temperatures up to at least 250°C. It, however, reacts with large numbers of gases, such as oxygen, nitrogen, hydrogen, carbon, carbon monoxide, carbon dioxide and methane, at temperatures above 300°C. This ability makes tantalum ideal as a getter in vacuum systems and gas purification systems. Generally, tantalum possesses good resistance to most liquid metals. Liquid metals that attack tantalum include aluminum, magnesium, cadmium, and zinc. 7 Aluminum Alloys N. J. Henry Holroyd Luxfer Gas Cylinders, Riverside, California I. INTRODUCTION The aim of this chapter is twofold: first, to briefly describe the various modes of corrosion (time-dependent environment-induced degradation) suffered by aluminum and its alloys; second, to discuss and review the influence played by surface films during these corrosion processes. This aspect has been chosen because although its importance is consistently highlighted in the published texts on aluminum corrosion (1–4), little detailed discussion ensues. A good example is the recently published book Corrosion of Aluminum Alloys (4) which reviews many aspects of the corrosion of aluminum and its alloys. The detailed discussion on the role of surface films during corrosion is limited to the statement, ‘‘Aluminum owes its excellent corrosion resistance and its usage as one of the primary metals of commerce to the barrier oxide film that is bonded strongly to its surface’’ and to a page of text giving basic information on aluminum films. Hopefully, this chapter will contribute to addressing this position. II. BACKGROUND Aluminum and its alloys are currently used in a wide range of applications, including aerospace, automotive, building, electrical, marine, packaging, and transportation. A comprehensive summary of the specific alloys (wrought and cast alloys) used for different applications has recently been published (4). Reasons why a particular alloy is selected are application dependent. For instance, for structural appli173 174 Holroyd cations, a potential user requires an alloy to provide an attractive strength-to-weight ratio, the ability to be joined using welding, adhesive bonding, and so forth and an acceptable corrosion resistance, all to be available at a competitive cost. For a nonstructural application such as a lithographic sheet, the user’s requirements are more focused on the uniformity of the alloy’s surface properties and their subsequent response to chemical and electrochemical processes. There are eight main groups of aluminum alloys: Alloy series 1000 2000 3000 4000 5000 6000 7000 8000 Main alloy groups 99.0–99.99% aluminum Al–Cu Al–Mn Al–Si Al–Mg Al–Mg–Si Al–Zn–Mg(–Cu) Others, some including Al–Li In most commercial applications the corrosion of aluminum and its alloys is not an issue during service life because the thin surface ‘oxide’ film formed during heat-treatment and\or during product fabrication reduces corrosion to negligible rates. In benign service environments this remains the case even if these preexisting surface films are mechanically disrupted because another protective film rapidly forms locally preventing the initiation of corrosion. Situations do arise, however, where this is not the case, and a surface that was protective, locally loses this capability and a corrosion process initiates. III. TYPICAL MODES OF CORROSION EXPERIENCED BY ALUMINIUM AND ITS ALLOYS Over the years, aluminum alloys have been reported susceptible to many forms of corrosion, some of which are rare and specific to aluminum alloys (for instance, ‘‘snowflake corrosion’’ or ‘‘fingerprint corrosion’’) (3). The more common forms of corrosion potentially suffered by aluminum and its alloys are both alloy-system and alloy-temper sensitive. A. General Corrosion General corrosion occurs when a preexisting ‘‘surface film’’ is no longer protective. This usually occurs in either strongly acidic or alkaline conditions, with Aluminum Alloys 175 the initial surface film ‘‘dissolving’’ and the new ‘‘film’’ that develops being nonprotective. The extent of the damage induced depends on the specific chemical composition of the electrolyte, its temperature, and its volume. At one extreme, the alloy can completely dissolve; at the other, the corrosion product ‘‘consumes’’ all the free water, resulting in a ‘‘gel’’ layer forming that can stifle further attack and eventually may become a ‘‘glasslike’’ solid layer. In highly acidic solutions (pH ⬍ 2), general corrosion usually occurs via a crystallographic mode, whereas in highly alkaline solutions (pH ⬎ 11), attack is usually associated with the formation of shallow ‘‘cusps.’’ General corrosion is not usually experienced during a commercial product’s service life unless the environmental conditions encountered stray well outside those anticipated for its application. B. Localized Corrosion Localized corrosion is the most common form of corrosion experienced by aluminum alloys during service applications. It occurs when the surface ‘‘oxide’’ film, for some reason, locally can no longer prevent the initiation of corrosion. The most common forms of localized corrosion are pitting, intergranular corrosion, crevice corrosion, galvanic corrosion, exfoliation, stress corrosion, and corrosion fatigue. Less common forms, although in some cases commercially important forms, include filiform corrosion, erosion–corrosion, poultice corrosion, biological corrosion, snowflake corrosion, and fingerprint corrosion. C. Pitting Corrosion Pitting corrosion is the most common form of localized corrosion suffered by aluminum and its alloys, and although numerous pitting studies have been undertaken since the 1930s, the universal agreement on how pits initiate is still awaited. An example of the pitting of an aluminum alloy is shown in Fig. 1 and a polished metallographic cross section through a typical pit is shown in Fig. 2. D. Intergranular Corrosion Intergranular corrosion (IGC), second to pitting corrosion, is probably the next most common form of corrosion suffered by aluminum alloys. A typical example of intergranular corrosion is given in Fig. 3. It is potentially more damaging than pitting because for a given loss of metal, the percentage reduction in a structure’s load-bearing capability is significantly higher for IGC and the observed tendency for self-stifling is less common than is found for pitting corrosion. Numerous studies have been published over the years, and compared to pitting corrosion, there is a greater consensus on the mechanisms involved; albeit, details may be alloy system and temper dependent. In general, underaged tempers 176 Holroyd Fig. 1 Pitting morphology before and after cleaning an aluminum alloy after long-term exposure to a saline environment (reduction 3⫻). Fig. 2 A polished metallurgical cross section through a typical pit shown in Fig. 1 (magnification 200⫻). Aluminum Alloys 177 Fig. 3 Typical intergranular corrosion in an aluminum alloy (magnification 200⫻). are more susceptible to intergranular corrosion that either peak or overaged tempers. E. Crevice Corrosion An opportunity for crevice corrosion occurs when the local geometry provides a region where a high surface area of the metal can be exposed to a relatively small volume of an aqueous solution that has a restricted access to the bulk environment. It has been argued that crevice corrosion is a variant of pitting corrosion, where the initiation phase of pitting is eliminated by the crevice itself, providing the restricted geometry (i.e., a very large pit) needed for the development of the local environmental conditions for corrosion initiation. Although this is mechanistically incorrect, it is not an unreasonable visualization of why crevices should be avoided in good designs and how tight crevices could lead to the initiation of localized corrosion. F. Galvanic Corrosion Galvanic corrosion of aluminum alloys usually only occurs as a result of poor design or incorrect material selection (Fig. 4). The observed form of attack depends on the actual environmental conditions, the aluminum alloy system, its temper, and the magnitude of the galvanic stimulation. Galvanic corrosion generally manifests itself as one of the other forms of localized corrosion, usually 178 Holroyd Fig. 4 Several localized corrosion modes occurring simultaneously: pitting, galvanic corrosion (mixed metal), and erosion–corrosion (sometimes referred to as impingement). either pitting or crevice corrosion; however, severe galvanic stimulation can lead to general corrosion. G. Exfoliation Corrosion Exfoliation corrosion is sometimes referred to as ‘‘lamella’’ or ‘‘layer’’ corrosion because, in cross section, it has the visual appearance of ‘‘flaky pastry’’ (Fig. 5). It usually only occurs in highly worked wrought alloys and the phenomena may be Aluminum Alloys 179 Fig. 5 Section of an agricultural vehicle showing severe exfoliation corrosion: magnification (a) 3/4⫻; (b) 15⫻. 180 Holroyd regarded as corrosion occurring simultaneously along multiple parallel paths associated with a given microstructural region adjacent to the alloys external surface. Exfoliation corrosion has been observed in both heat-treatable and nonheat-treatable alloys. For heat-treatable alloys, the corrosion paths follow grain boundaries and/or regions adjacent to them, whereas for non-heat-treatable alloys, the regions followed involve those rich in transitional metals and the regions immediately adjacent. H. Stress-Corrosion Cracking Stress-corrosion cracking (SCC) is a complex phenomena requiring the concurrent action of both local strain and corrosion. Generally, it is only of practical concern for the higher-strength aluminum alloys, and then only when heat-treatable alloys [e.g., Al–Cu–Mg, Al–Zn–Mg(–Cu), or Al–Li] are in underaged or peak-aged tempers. Non-heat-treatable Al–Mg alloys with magnesium concentrations above around 3% (wt) may be susceptible to SCC if they become sensitized by a grain-boundary precipitation of an active phase during service life. (This process tends to be accelerated in highly worked alloy microstructures.) Medium-strength heat-treatable alloys such as Al–Mg–Si (6xxx series) alloys are rarely susceptible to SCC other than when highly alloyed and subjected to unusual noncommercial laboratory heat treatments (3,5). Several extensive and complete reviews of the SCC of aluminum alloys have been published (5–9) and the reader is directed to these works, as this aspect of aluminum corrosion will not be covered in this chapter. I. Corrosion Fatigue Corrosion fatigue (CF) is a complex phenomenon requiring the concurrent action of both local cyclic strain and corrosion. CF can occur in a wide spectrum of conditions, involving crack-propagation mechanisms ranging from where local corrosion is contributing to a process that is essentially pure fatigue through to a situation in which the cyclic loading is enhancing the crack growth rate in a stress-corrosion process. As with SCC, the reader is directed to an extensive and complete review of the subject (10). IV. ENVIRONMENTAL EFFECTS ON ALUMINUM ALLOYS Environmental effects on aluminum and its alloys are dictated initially by how the surface ‘‘oxide’’ responds to its new environment. Before discussing these interactions, it is appropriate to review the types of surface ‘‘oxide’’ film that typically are present on the aluminum surfaces prior to their environmental exposure. Aluminum Alloys 181 A. Surface ‘‘Oxide’’ Films on Aluminum Aluminum is an active metal and its stable behavior depends on its exposed surfaces being covered by a protective thin ‘‘oxide’’ film. The nature of this film, its stability over time, and the consequences following its local disruption in a given set of environmental conditions usually dictates the corrosion performance of aluminum-based alloys in those conditions. Despite this knowledge, it is interesting to note how little attention has been paid to this detail over the years by the numerous research workers studying the corrosion and electrochemistry of aluminum alloys. The first reaction is to remove the ‘‘irreproducible’’ as-received surface and replace it with something else. Good examples of where the corrosion performance of a given alloy in a particular environment is strongly influenced by the nature of the pre-existing surface ‘‘oxide’’ prior to the alloy being exposed to an environment include the following: 1. Godard’s (11) statement: ‘‘Experiments I have performed indicate that the thickened films developed in pure water at room temperature increase the resistance of the surface to corrosion; if such films can be developed before some corrosion conditions are encountered no corrosion will occur, whereas a freshly exposed surface with only an air formed film will be subject to corrosion.’’ 2. The ‘‘blue’’ corrosion phenomenon (12), experienced in the early 1980s, where users of household aluminum foil (⬃99% Al) observed small ‘‘blue’’ patches developing on foil surfaces. This corrosion issue occurred because the foil alloy contained unexpected trace levels of lithium (⬍10 ppm) that, during thermal treatment, resulted in the oxide film having localized areas of lithium surface enrichment of up to 4000: 1 (13,14) which upon exposure to humid air (as often found in domestic kitchens) suffered corrosion and generated a local ‘‘blue’’ color, whereas regions without lithium suffer no attack. 3. Ranking of the pitting corrosion susceptibilities of aluminum alloy systems using well-known standard methods can be influenced by the surface pretreatment employed prior to testing (15). B. Oxide Films Formed on Aluminum Alloys in ‘‘Dry’’ Environments The surface films forming on aluminum exposed to the natural atmosphere, dry air, or oxygen at room temperature initiate instantaneously. These films consist ˚ thick) covered by a slightly thicker, less of a compact inner layer (10–15 A dense outer layer of an amorphous noncrystalline aluminum oxide that reaches ˚ after a few hours (16). a limiting thickness, typically less that 50 A These duplex ‘‘oxide’’ films are produced on all aluminum alloys for tem- 182 Holroyd peratures up to around 400°C, as long as the alloy does not contain alloying elements, such as magnesium or lithium, that react selectively to form their own oxides. The principle difference for these higher-temperature films is that the ˚ -thick compact thickness of the outer γ-Al2O3 layer, forming over the 10–15-A inner layer, increases with temperature. Film thickness measurements quoted in the literature vary significantly, reflecting that, as with corrosion, the surface preparation prior to oxidation often has a significant influence on the subsequent oxidation behavior (17). Notwithstanding this, it is reasonable to assume oxides will ˚ thick [normally ⬍ 100 A ˚ (18)] for temperatures up to at 400°C be less than 200 A ˚ and increase to around 400 A at 600°C (19). Alloying elements such as copper and manganese have been reported to enter and dope γ-Al2O3 at these higher temperatures (20). At temperatures above about 425–450°C, the amorphous outer layer of the duplex film initially forming on aluminum alloys undergoes a discontinuous structure change. This allows the rapid migration of oxygen to the oxide–metal surface and the nucleation and growth of crystalline γ-Al2O3 as a new phase below the amorphous layer (21,22). The presence of alloying elements that react selectively, forming their own oxides, slightly complicates the oxidation of aluminum alloys. For commercial aluminum alloys, the principal elements doing this are magnesium and lithium. For aluminum alloys containing more than ⬃0.5 wt% magnesium, the oxides generated at temperatures up to around 300°C are independent of the magnesium addition. However, for higher temperatures, the outermost surface of the oxides formed is essentially a dense layer of MgO crystals (23). Electron microscopy (24) has shown that MgO crystals can also form at the metal–oxide interface, thereby disrupting the inner compact alumina layer. Zinc additions to aluminum alloys appear to have minor effects on an alloy’s oxidation characteristics (25,26) with no evidence existing for zinc incorporation into the films formed on high-purity Al–Zn–Mg alloys (27) and only limited evidence for commercial Al–Zn–Mg–Cu alloys (26,28). Experimental evidence for these alloys shows that the alloy grain boundaries provide a shortcircuit diffusion path for magnesium to participate in the oxidation process (27). The presence of lithium in aluminum alloys even at low concentrations has a dominant effect on the oxidation behavior of aluminum alloys, overriding those due to magnesium and other alloy additions to commercial alloys (25). C. Oxide Films Formed on Aluminum Alloys in ‘‘Wet’’ Environments As stated earlier, the natural oxide formed on aluminum in ‘dry’ air at room ˚ -thick barrier layer next to the temperature is duplex with a compact 10–15-A metal covered with a thicker less dense layer. The thickness of the outer layer, although initially around 50 angstroms, is highly sensitive to the prevailing envi- Aluminum Alloys 183 ronmental conditions (11,29). Although researchers have studied the filming characteristics of aluminum alloys in water for decades, accurate prediction of the expected film thickness of aluminum alloy is not yet possible. This is due to the multiplicity of variables involved which include: the nature of the pre-existing surface oxide, details of the water chemistry composition, temperature, pH, oxygen content, etc. and the alloy composition. An excellent review of the films forming in the aluminum-water system has been provided by Alwitt (29), who states: ‘‘It is an interesting fact that those phases commonly present as surface films on aluminum—pseudo-boehmite, bayerite, boehmite and γ-Al2O3 —are all metastable phases of the Al2O3 –water system’’ and ‘‘the film composition is often the most kinetically accessible product, which often is only an intermediate product on the path to thermodynamic equilibrium.’’ Film forming on aluminum exposed to water at various temperatures can be categorized by four temperature regimes: a) room temperature up to around 40°C; b) above 40°C to 100°C; c) 100°C to 150°C and d) above 150°C. 1. Films Formed at Temperatures at and Below 40°C These films differ from those formed at temperatures up to 100°C (29–31), with the films formed at 40°C being basically either exclusively bayerite or a very thin of layer pseudo-boehmite covered by a thicker layer of bayerite. An explanation of this behavior at temperatures below around 45°C could be that although pseudo-boehmite can form on aluminum surfaces immersed in water at these temperatures, its growth rate is so slow and bayerite precipitation occurs because it is kinetically more favorable. Alternatively, the pseudo-boehmite could start to form, but it then dissolves at the film–solution interface and reprecipitates as the thermodynamically more stable product bayerite (29,30), thereby generating a duplex film with a very thin partially protective pseudo-boehmite layer covered by a thick bayerite layer. The net effect is that the films formed at temperatures below 45°C are thicker, more porous, less protective, and more brittle, with a tendency to locally spall away from the metal substrate (29,30) (e.g., see Fig. 12 in Ref. 29). A clear indication that the filming characteristics change at temperatures around 45°C is provided by the temperature dependence of the film formation induction times (Fig. 6). The implied activation energy is 18.7 kcal/mol for temperatures in the range 50–100°C, whereas it is significantly higher for lower temperatures. 2. Films Formed at Temperatures Above 40 °C and Below 100 °C These films have a duplex structure, consisting of an inner layer of pseudoboehmite [a poorly crystallized oxyhydroxide, similar to boehmite (AlOOH) but 184 Holroyd Fig. 6 Temperature dependence of the induction time for film growth on pure aluminium exposed to distilled water in the temperature range 30–100°C showing that filming process change occurs at temperatures around 45°C. 䊊 from Ref. 18; 䊉 from Ref. 29; 䉱 from Ref. 89; and ⫻ from Ref. 33. containing excess water (18)] and an outer layer of bayerite crystals [a form of aluminum hydroxide (32)]. Film growth for this temperature range proceeds in three stages: (1) an induction period, during which no significant weight change is observed, (2) a period during which pseudo-boehmite forms, and (3) bayerite crystallization onto the pseudo-boehmite layer. Stages 2 and 3 can overlap, with bayerite crystal nucleation occurring on the boehmite while the boehmite is still growing (29). It has been suggested that the cessation of film growth is dictated by the completion of the bayerite layer and is independent on the formation of the underlying pseudo-boehmite layer, which is considered to reach a limiting thickness rapidly (29). The precise mechanism is not fully established. Scamans and Rehal (32) have made a reasonable proposal based on published work, suggesting that the pre-existing amorphous film thickens during the induction period [as proposed Aluminum Alloys 185 by Hart (33)] with either proton, hydroxyl ions, and/or water penetrating the film to establish a soluble aluminum species. A dissolution–precipitation process follows, leading to the rapid formation of a pseudo-boehmite layer on the aluminum surface that, once complete, can only thicken slowly via a process limited by the diffusion of water into and through the film, as opposed to the outward diffusion of Al3⫹ ions. These authors suggest that the mechanistic details of the bayerite formation process awaits clarification. However, it probably involves a further dissolution–precipitation process, in this case, occurring at the pseudoboehmite–solution interface with a typical bayerite layer thickness growing to 2–3 µm at 70°C. 3. Films Formed at Temperatures in the Range 100–150 °C These films on aluminum in water and steam have been studied in great detail and reviewed by Altenpohl (34). According to Altenpohl (35,36), typical boehmite films developed on superpurity aluminum immersed in boiling distilled water involve two types of boehmite. The first forms within 10 min in boiling water and will redissolve fairly readily, whereas the second, which is highly insoluble, forms subsequently as a thinner layer beneath the thicker, more soluble form of boehmite. A schematic of these boehmite films is given in Fig. 7 and typical film Fig. 7 Schematic representations as given by Altenpohl (34) for cross sections through boehmite films formed on superpure aluminum immersed in boiling distilled water for 1 ˚ -thick barrier layer at h and 5 h. [Note: Subsequent studies indicate that a retained 10-A the aluminum–film interface is unlikely to be present (29).] 186 Holroyd Fig. 8 Typical film thickness–time data according to Altenpohl (34) for film growth on AA1145 immersed in boiling distilled water. thickness–time data for their film growth on AA1145 exposed to distilled water is shown in Fig. 8. It is worth noting that Altenpohl’s filming model (34) envisions that the thin, compact ‘‘barrier’’ oxide layer on the aluminum surface prior to any reaction with water remains in situ throughout the boehmite filming process. Alwitt (29), in his review of the aluminum–water system states that in his opinion, ‘‘there appears to little experimental evidence for such a film.’’ The benefits offered by ‘‘boehmite’’ films can be impressive. For instance, data from Leidheisser and Kellerman’s study of straining aluminum alloy wires in aqueous environments (37) shows that immersing commercially pure aluminum (AA1100) in boiling water for 15 min prior to rapidly straining in a 0.1M NaCl solution significantly improves the surface films ductility compared to that given by an air-formed film or an anodized film exposed to boiling water (Fig. 9). Another good example is the improved pitting-resistant ‘‘boehmite’’ films can provide in potentially aggressive tap waters, as is shown in Table 1. Based on Aluminum Alloys 187 Fig. 9 Effect of surface films on the electrochemical potential response after commercially pure aluminum (AA1100) wires are rapidly plastically strained 1.5% in a deaerated 0.1M NaCl solution, pH 5.5. (Data from Ref. 37.) the data in Table 1 and other data, a significant interest was generated during the 1960s to attempt to utilize ‘‘boehmite’’ films commercially to render aluminum cans suitable for packaging beer, milk, and other products (34). Attractive as the ‘‘boehmite’’ films are, their commercial exploitation to date has been restricted by issues associated with the conditions needed to successfully generate these films in either boiling water or in high-temperature steam. For instance, 1. Protective ‘‘boehmite’’ films have only commercially been reproducibly formed on high-purity aluminum. 2. Trace additions of contaminants in boiling water or steam can either stifle film growth or promote corrosion and/or the growth of a nonprotective film. Re-examination of the boehmite film thickness data provided by Altenpohl (34) for six superpure aluminum alloy compositions subjected to different surface pretreatments prior to their immersion in boiling high-purity distilled water for 188 Holroyd Table 1 Number of Pits Generated on ‘‘Filmed’’ Commercially Pure Aluminum (AA1145) After Exposure to Fairly Hard Tap Water for 9.5 Months at Room Temperature No. of corrosion pits per dm2 in tap water Type of distilled water used for boehmite film formation 50°C 70°C A B C A ⫹ 0.1N NH4OH 5 30 12 0 2 3 0 0 B ⫹ 0.1N NH4OH C ⫹ 0.1N NH4OH Air-formed film 0 0 18 0 0 3 Specific resistance of distilled water (Ωcm) 8 ⫻ 10⫺5 7 ⫻ 10⫺5 5 ⫻ 10⫺5 2.18 ⫻ 10⫺4 to 1.45 ⫻ 10⫺4 (start) (finish) Note: Boehmite films produced by immersion in boiling distilled waters for 4 hours, with and without a 0.1N ammonium hydroxide addition. Source: Data from Ref. 34; see reference for details on distilled-water compositions. various times reveals that the film thickness generated is effectively controlled by an alloy’s iron, copper, and silicon contents. This conclusion is based on the following observation of high-purity aluminum (i.e., concentrations of iron, copper, and silicon will always be low): 1. Thicker films only occur if an alloy’s copper and iron contents are below 0.0015 and 0.002 wt%, respectively (Fig. 10). 2. When copper and iron levels are sufficiently low, increasing an alloy’s silicon weight percent favors thicker films (Fig. 10). The latter observation is interesting in view of the well-known detrimental effects trace additions of silicon to the alloy have on film formation at temperatures above 150°C or dissolved silicate ions additions to boiling distilled water have on boehmite film formation (Fig. 11). Based on the latter, it is reasonable to suggest that it is the ‘‘free silicon’’ in the alloy that is detrimental to boehmite film formation, and when at appropriate low concentrations, iron and silicon can form relatively benign precipitates, thereby allowing copper levels to control boehmite film thickness (see Fig. 10). Some published data on the influence of anions on the growth behavior of boehmite films on aluminum immersed in boiling water are available (34,38,39). Data presented by McCune et al. (38) for 99.99% aluminum (AA1100) in boiling Aluminum Alloys 189 Fig. 10 Effect of silicon content on the boehmite film thickness developed on a range of superpurity aluminum compositions immersed in boiling distilled water for either 2 or 8 days. (Data from Ref. 34.) 190 Holroyd Fig. 11 Boehmite film thickness developed on 99.99% aluminum after 4 h immersion in boiling distilled water containing various silicate ions additions up to 15 ppm. (Data from Ref. 34.) distilled and deionized water containing a range of inhibitor anions as single additions suggests that in all cases the film growth was less than the pure water case after 1 h (Fig. 12). Comparison of these data with that provided by Altenpohl (34) indicates that although there is good agreement between the experimental data, the addition of anions does not necessarily lead to film thickness reductions. 4. Films Formed at Temperatures Above 150 °C At atmospheric pressure, boehmite is the only phase consistently reported (40– 43) to form on aluminum surfaces exposed to water or steam in the temperature range 150–374°C, the critical temperature of water. Hart and Maurin (41) have observed that Dispore can form as a thin outer layer over the boehmite after extended periods at temperatures above 300°C. For temperatures above 374°C, γ-alumina and corundum are formed (41). In the mid-1950s, there was considerable interest in using aluminum to clad the fuel elements used in water-cooled nuclear power reactors, with the Aluminum Alloys 191 Fig. 12 Apparent pseudo-boehmite film thickness formed on 99.99% (AA1100) aluminum immersed in boiling distilled water containing various single-anion additions as sodium salts (1 g/L nitrate, 2 g/L sulfate, and 1 g/L silicate) for various up to 1 h. (Data from Ref. 38.) aluminum being exposed to water at temperatures up to around 350°C. Aluminum’s attractions for this application were its low cost, ease of fabrication, and low-adsorption cross section for neutrons. An intense research and development (R&D) effort ensued in North America to overcome two issues: 1. To improve aluminum’s environmental performance in water at elevated temperatures and to avoid the situation where, above a critical 192 Holroyd temperature, surface films lose their protective nature and allow intergranular corrosion and rapid loss of metal 2. To increase the strength of aluminum at temperatures around 300°C Aluminum alloy development was the R&D approach adopted to achieve the above. This decision was based on the belief that appropriate alloying additions would exist that could simultaneously increase an alloy’s strength while also increasing the critical temperature above which aluminum corrodes rapidly, which was known to increase as the alloy’s purity decreased (44). [Typical critical temperatures are 105°C for 99.99% aluminum, 150°C for 99.85% aluminum, 180°C for 99.5% aluminum and a higher temperature for 2S aluminum (99.0–99.3% aluminum), which normally performs well in water up to temperature around 200°C (34).] Work in the United States at the Argonne National Laboratories in 1955 established that a 1 wt% nickel addition to AA1100 prevented intergranular attack at 350°C (45), changing the mode to uniform corrosion (40). This led to the development of the experimental alloy X-8001, which was superseded in 1957 by two further alloys based on Al–Ni–Fe–Ti (A203X and A198X), that both offered a superior corrosion performance in water at 360°C, which was attributed to these alloys having a finer dispersion of second-phase precipitates. Corrosion performances for the initial commercially cast versions of the above Al–Ni–Fe–Ti alloys in high-temperature water were extremely variable, and after detailed studies, it became evident that the corrosion performance was extremely sensitive to the alloy’s purity level—in particular, the silicon weight percent. A further alloy development program was undertaken by the General Electric Company to generate an alloy that provided yet a better corrosion performance in water at 360°C while also having a sufficient tolerance to the alloy’s purity level that it could be readily cast by conventional casting practices. An alloy based on Al–1.8 Fe–1.2 Ni resulted (46) that, as well as being sufficiently tolerant to the alloy’s purity levels, was also tolerant of long-term thermal treatment at 550°C and provided a superior corrosion performance. In parallel with the United States effort, Krenz (42), working in Canada, reported two Al–Ni–Fe–Si alloys (Al–2 Ni and Al–0.5 Ni, both containing 0.5% Fe and 0.2% Si) that performed reasonably well when exposed to water at 300°C. Initial attack of the alloys was rapid, but further attack slowed down to a linear rate that was marginally lower for the higher-nickel-containing alloy. MacLennan (47) exposed the above two Al–Ni–Fe–Si alloys along with 2S aluminum to water at 300°C for short periods of time (minutes) and then took replicas of the inner and outer surfaces of the oxide films and evaluated them using electron microscopy. Based on this study and further work involving longer exposure times in water at temperatures in the range 150–340°C (48), MacLennan attributed the improved corrosion performance for the nickel-containing alloys to their oxide films providing a local ‘‘stress relief effect.’’ This stress relief effect was deemed to minimize any tendency for the films to crack, as is Aluminum Alloys 193 thought to occur to the oxides formed on aluminum exposed to high-temperature water (49). Improved film ‘‘ductility’’ was suggested to be provided by the second-phase particles and the adjacent aluminum-forming local cells where the aluminum rapidly dissolves until the particle is isolated from the metal. Other workers in Canada simultaneously evaluated the oxide films forming on these Al–Ni–Fe–Si alloys after exposure to pure water at high temperatures (200–300°C). Grenblatt and McMillan (50), using electrochemical polarization techniques, established that the oxide films differed from those formed on AA1100, and Greenblatt (51), using optical metallography to study the oxide thickness and the corrosion pit depths, showed that the corrosion mechanism for AA1100 differed from that for Al–Ni–Fe–Si alloys. Early operational experience for aluminum-alloy-clad fuel elements used in pressurized and boiling-water nuclear reactors quickly revealed that the corrosion performance was significantly inferior in the pressurized-water reactor case with the aluminum alloy’s oxide film ‘‘dissolving’’ in the high-temperature water at appreciable rates (52). The reason for this is that in the pressurized-water case, unlike for the boiling-water case, the oxide surface temperature is not uniform and the solution layer in contact with the oxide is not always nearly saturated with dissolved oxide. 5. Electrochemical Potential–pH Diagrams Electrochemical potential versus solution pH diagrams for the aluminum–water system have been presented for temperatures in the range 25–300°C (53,54). Construction of these diagrams are based on calculations using thermodynamic data which inevitably involves making assumptions that are known to introduce oversimplifications. For instance, for the calculations, it is normally assumed that the chosen stable oxide film is in equilibrium with soluble aqueous species such as Al3⫹ and Al(OH)4⫺ ions, although it is well established that at intermediate pH’s, the hydrolysis products of the Al3⫹ ion dominate, and at higher temperatures, polynuclear complexes readily form. Despite these oversimplifications, qualitatively useful information may be gleaned from these diagrams. For example, (a) the pH range for oxide stability is highly temperature dependent [4–10 at 25°C, 3–7 at 100°C, and 1.5–3.5 at 300°C (54)] and (b) the solution pH for minimum corrosion as a function of temperature are correctly predicted [⬃6.5 at 50°C (40), ⬃5.7 at 92°C (55), and ⬃3.0 for 170–300°C (40)]. A typical potential–pH diagram for aluminum in water at 25°C is given in Fig. 13. V. OXIDE FILM DEVELOPMENT AND CORROSION INITIATION IN AQUEOUS ENVIRONMENTS Aluminum surfaces are normally covered by a protective thin ‘‘oxide’’ film. The specific nature of this film depends on the conditions under which it was formed 194 Holroyd Fig. 13 Electrochemical potential–solution pH diagram (often referred to as a Pourbaix diagram) showing thermodynamically predicted regions for corrosion, passivation, and immunity of aluminum in water at 25°C, assuming that the oxide formed is bayerite. and subsequently stored. Time-dependent processes are initiated when these surfaces are exposed to aqueous environments, the results of which are dependent on several factors, including the aluminum alloy composition, temper and microstructure, the nature of the pre-existing surface ‘‘oxide’’ film itself, and the prevailing local electrochemical and environmental conditions (e.g., electrode potential, solution chemistry, temperature, pH, oxygen concentration, etc.). When exposed to sufficiently aggressive aqueous environments, the protection provided by the pre-existing ‘‘oxide’’ film on aluminum alloys is negated and general corrosion of the underlying metal is initiated. For the vast majority of structural or engineering applications of aluminum alloys, these situations are of no interest and these conditions need to be avoided at all costs. Exceptions to this usually involve situations in which the alloy is intentionally being used as (a) a battery anode [e.g., an aluminum–air battery (56)] or (b) a sacrificial anode (57) or aluminum-based spray coating (58–60) to protect another material Aluminum Alloys 195 such as a steel tension leg in an offshore oil platform structure (61). In these cases, the general corrosion behavior itself has to be ‘‘controlled’’ to maximize coulombic efficiencies (e.g., by alloy composition and microstructure manipulation). Another exception is the situation in which an aluminum alloy is subjected to a short period of alkaline etching aimed at removing the pre-existing surface film and the preferential removal of second-phase particles that otherwise would act as a cathodic reaction sites during any subsequent corrosion processes (15). Although it is true that aluminum alloys are only resistant to general corrosion when the solutions’ pH’s are neither too high nor too low, it is now generally accepted that an alloy’s corrosion performance is controlled by the nature of the aluminum salt involved rather than the solution’s pH produced by the hydrolysis of the aluminum cation (62). Published information on the inhibition of aluminum alloy corrosion in highly acidic (63–65) and highly alkaline (66–68) environments is available. To simplify the following discussion we will initially concentrate on the behavior of nonalloyed aluminum (i.e., high-purity and commercial purity grades of aluminum alloys). Following this, we will examine the effects of single-alloy additions (i.e., binary alloys) and, finally, consider the commercial alloy systems. A. Unalloyed Aluminum Alloys Although some controversy still exists in the literature on the details of corrosion initiation, there is now a strong opinion that the initial process occurring when an ‘‘oxide’’-filmed aluminum surface is first introduced into an aqueous environment, although being sensitive to electrochemical potential, is predominantly a ‘‘chemical’’ process. In a comprehensive review covering over 70 years of published literature Foley (69) came to the conclusion that the initiation of localized corrosion of aluminum could be described by four steps involving the following: 1. Adsorption of a reactive species (usually an anion) on the ‘‘oxide’’covered aluminum alloy surface 2. A chemical reaction of the adsorbed species with the aluminum ion in the aluminum oxide lattice or precipitated aluminum hydroxide 3. The effective ‘‘thinning’’ of the ‘‘oxide’’ film by ‘‘dissolution’’ or anions penetrating the film 4. ‘‘Direct’’ attack of the exposed metal by the anion, possibly assisted by anodic potentials (sometimes, Steps 3 and 4 can occur simultaneously) Step 1: Adsorption A range of analytical tools, including autoradiography, secondary ion mass spectroscopy (SIMS), and X-ray photoelectron microscopy (XPS), have provided substantial experimental evidence for the competitive adsorption of anions on aluminum surfaces [e.g., chloride (70–75), sulfate (38,70,72,75), perchlorate (72), nitrate (38,72), 196 Holroyd chromate (72,76), and molybdate (38,77,78)] and the subsequent influence upon whether or not aluminum will subsequently undergo dissolution. This is particularly the case for the adsorption of chloride ions as a preliminary step to aluminum alloys suffering pitting corrosion (70– 75). Observed adsorption behavior is dependent on the anion species involved. For example, experimental evidence provided by Tomcsanyi et al. (70) using a radio-tracer technique on 99.99% pure aluminum at its free-corrosion potential in various aqueous solutions has shown the following: 1. Adsorbed labeled sulfate ions can be displaced by further additions of nonlabeled sulfate ions, whereas they are not displaced by the further additions of chloride ions (Fig. 14). 2. Chromate ions displace adsorbed chloride ions and electrochemical potentials move to more positive values (Fig. 15), whereas further additions of nonlabeled chloride ions do not displace adsorbed labeled chloride ions. Fig. 14 Accumulation of labeled sulfate ions on pure aluminum surfaces: (A) exchange when additional nonlabeled sulfate ions are added and (B) effect of adding 5 ⫻ 10⫺3M sodium chloride. (Data from Ref. 70.) Aluminum Alloys 197 Fig. 15 Effect of chromate ion additions to the surface accumulation of chloride ions on pure aluminum in a 1M sodium sulfate solution: (A) a 0.1M Na2CrO4 addition to a solution containing 4 ⫻ 10⫺2M sodium chloride and (B) a 5 ⫻ 10⫺2M Na2CrO4 addition to a solution containing 2 ⫻ 10⫺2M sodium chloride. (Data from Ref. 70.) 3. The concentration of adsorbed chloride is both time and concentration dependent (Fig. 16). 4. The excess surface concentration of chloride ions is insensitive for all electrode potentials up to around ⫹750 mV on the saturated calomel electrode scale (SCE) other than for cathodic potentials where measurable hydrogen evolution occurs and adsorbed concentrations significantly decrease with decreasing potential. 5. Surface concentrations of adsorbed chloride ions are not fully reversible when the electrode potential is moved from a given anodic value to a sufficiently cathodic one and then back to the original anodic potential (Fig. 17). These results are important. The first two give a good indication of how chromate and sulfate ions initially inhibit the detrimental effects of chloride ions [note: Augustynski (72) suggests that nitrate ions retard the 198 Holroyd Fig. 16 Time dependency of the surface accumulation of chloride ions on pure aluminum exposed to various bulk solution chloride concentrations in a 1M sodium sulfate supporting electrolyte. (Data from Ref. 70.) adsorption of chloride ions], whereas the next two are consistent with the adsorbed chloride ions ‘‘chemically’’ rather than ‘‘electrochemically’’ reacting with the ‘‘oxide’’ film. Justification of the latter suggestion is that the adsorption behavior for a given surface condition is insensitive to electrochemical potential and that although the concentration of the adsorbed chloride increases with increasing bulk chloride solution, the absolute concentrations always start decreasing after ⬃1 h. Rationalization of the final observation that the adsorption behavior was not reversible was that the adsorption is sensitive to the surface ‘‘roughness’’ of the aluminum substrate, which will be modified by the electrode potential being held at a cathodic potential. Tomcsanyi et al.’s (70) suggestion that chloride adsorption on oxide-covered aluminum is insensitive to electrode potential would seem inconsistent with Augustynski’s (72) results showing that adsorbed chloride concentrations increased from around 3 at% at the free-corrosion potential to around 12–13 at% at potentials approaching the critical pitting potential. Aluminum Alloys 199 Fig. 17 Evidence that the concentration of accumulated chloride ions on aluminum surfaces depend on both time and the current surface condition. Effects of cathodic polarization were not reversible. (Data from Ref. 70.) Rationalization of these differences is probably associated with the timedependent changes occurring after the chloride ions are initially adsorbed; this will be discussed below in Step 2 of the proposed process, which covers the chemical reaction of the adsorbed species with the substrate. Further evidence supporting the proposal that chloride ions adsorb on aluminum surfaces and then react with the ‘‘oxide’’ is provided by the simple but elegant experimental work published by Nguyen and Foley (73), who measured the uptake of chloride ions and the production of Al3⫹ ions as a function of chloride concentration when 10 g of either aluminum powder or aluminum oxide were mixed with 100 mL of an aqueous sodium chloride solution for 16 h. The results generated showed that the chloride ions adsorption is consistently higher on aluminum oxide than on freshly generated aluminum powder, and for both substrates, the adsorption increased with increasing bulk solution chloride concentrations up to a maximum at around 0.8M sodium chloride. The results for alumina are summarized in Fig. 18, along with data from similar experiments on aluminum powder conducted by Drazic et al. (74), who used 200 Holroyd Fig. 18 Adsorption of chloride ions onto alumina from a range of sodium chloride solutions. ⫻ ⫽ from Ref. 73; 䊉 ⫽ from Ref. 74. more dilute sodium chloride solutions. (Drazic et al.’s aluminum powder behaved as if it was alumina, so it was assumed it was coated with a relatively thick layer of aluminum oxide.) Having established that anion adsorption on oxide-covered aluminum immersed in aqueous environments is a competitive process, with certain ions promoting passivity and others film degradation, the logical question to ask is whether anion adsorption is a uniform or localized phenomenon. An argument for anion adsorption on alumina being relatively uniform can be made on the basis that the measured isoelectric point for alumina (79) and aluminum-oxide-coated aluminum (80) is in the range 8.9–9.1, irrespective of the chloride concentration (79). Hence, it is reasonable to expect that negatively charged species such as anions will be uniformly attracted to aluminum-oxide-coated aluminum surfaces. In line with this, Nguyen and Foley’s (73) results show that chloride ion adsorption on aluminum oxide exposed an aqueous sodium chloride solution of given Aluminum Alloys 201 concentration increased linearly on a one-for-one basis with the increasing surface area of alumina. An argument for anion adsorption being locally enhanced can be made using the fact that metal surfaces are rarely homogenous and enhanced localized adsorption and ‘‘surface activity’’ will tend to occur preferentially at imperfections, weak spots, and flaws in the surface oxide. In the early 1970s, Richardson and Wood (81,82) proposed that oxide ‘‘flaws’’ in aluminum oxide act as ‘‘active centers’’ for the initiation of localized corrosion. They supported their proposal with experimental evidence from transmission electron microscopy (TEM) studies for the existence of large ‘‘flaws’’ in aluminum oxide layers electrochemically grown on aluminum substrates. Further studies by Wood and later by Thompson and their co-workers have led to numerous publications generating evidence demonstrating that the initiation of the pitting of aluminum could be via an oxide ‘‘flaw’’ mechanism. Irrefutable evidence is awaited; however, the recent evidence from ultramicrotomed sections of pure aluminum exposed to an aqueous environment (83) is very convincing. Step 2: Chemical Reaction of the Adsorbed Species with the Surface ‘‘Oxide’’ Evidence for the solubility of oxide films on aluminum substrates immersed in aqueous environments was provided by Lorking and Mayne (84,85) in the early 1960s. These workers generated nonporous anodized layers on aluminum and then exposed them to a wide range of anions in aqueous environments which included chloride, sulfate, benzoate, acetate, citrate, phosphate, and chromate anions. In all cases, dissolved aluminum species were detected within 24 h and the concentrations detected were directly related to the initial corrosion rates for all the environments evaluated other than chloride. On the basis of these findings, the authors concluded that the initial anodic reaction in the chloride solutions was the formation of soluble aluminum chloride, whereas in all the other aqueous solutions, the anions promoted the repair of the anhydrous oxide. Perhaps the clearest evidence that a chemical reaction occurs between aluminum oxide and an aqueous environment is provided by the experiments conducted by Nguyen and Foley (73). These experiments involved the mixing of 10 g of aluminum oxide or 10 g of fresh aluminum powder with 100 mL of various aqueous saline solutions for 16 h and the subsequent measurement of the chloride ions adsorbed and the aluminum species generated. The data obtained indicated that while the chloride ion concentration adsorbed on the alumina increases with increasing initial chloride concentration (Fig. 18), the concentration of the dissolved aluminum species found in the aqueous solution decreases, with the highest concentration being recorded when the alumina is exposed to doubly- 202 Holroyd Fig. 19 Aluminum ion concentration (Al3⫹) resulting from immersing 10 g of alumina or fresh aluminum powder in various aqueous sodium chloride solutions. (Data from Ref. 73.) distilled water (Fig. 19). Results for the aluminum powder were different in that the dissolved aluminum concentrations were considerable lower and independent of the solution chloride concentration, and although the adsorbed chloride ions increased with increasing solution chloride concentration, the absolute concentrations were always considerable lower than on alumina (73). Based on the above results, it is reasonable to conclude that aluminum oxide in the presence of water undergoes hydration and that when chloride ions are available, the positive surface charge will lead to adsorption and the generation of aluminum hydroxy–chloro species such as Al(OH)2Cl and Al(OH)Cl2 or soluble complexes of the form Al(OH)2Cl2⫺ (86). Even more compelling evidence supporting the latter proposal was provided by Nguyen and Foley (73) when they mixed 10 g of alumina or aluminum powder with 100 mL in a 1N sodium chloride solution acidified with 0.01N aluminum chloride. In the experiments involving alu- Aluminum Alloys 203 mina, the solution’s pH became alkaline along with the dissolved aluminum concentration falling dramatically from its initial value of 83 mg/ L to a final value of 3.29 mg/L and the formation of a heavy gelatinous precipitate, presumably via a chemical reaction of the type Al3⫹ ⫹ 2OH⫺ ⫹ Cl⫺ → Al(OH)2Cl For the aluminum powder case, the dissolved aluminum concentration increased from 82.9 to 140.0 mg/L and the solution pH remained acidic. Presumably, in this case, the ‘‘dissolution’’ of the thin layer of ‘‘oxide’’ on the aluminum powder occurs without the solution’s pH increasing significantly and the surface layers of the aluminum powder can then ‘‘dissolve’’ in the acidic solution. The generation of a heavy gelatinous precipitate in the case when Nguyen and Foley (73) mixed alumina with an acidified saline solution provides clear evidence of the ‘‘chemical reactivity’’ of aluminum oxide in acidic aqueous environments. Another good indicator of the ‘‘chemical reactivity’’ of oxide films on aluminum surfaces exposed to aqueous saline environments is given by Tomcsanyi et al.’s (70) observation that the concentration of excess chloride adsorbed on aluminum surfaces decrease with time (Fig. 16). This phenomenon was deemed by Tomcsanyi et al. (70) to be a multistep heterogeneous transformation of the original air-formed oxide into a mixture of oxo-hydroxo- and chloro- complexes, with the details of the process being influenced by the initial surface oxide and the local solution chemistry at the oxide–solution interface. It could be wrongly concluded that chemical reactivity between the initial surface ‘‘oxide’’ and an aqueous environment necessitates the presence of aggressive species. Clear evidence that aggressive species are not essential is provided by evidence published by Scamans and Rehal (32). In their electron microscopy studies, small surface blisters were observed to form in the oxide when a wide range of aluminum-oxide-covered aluminum alloys (pure Al, Al–Mg, Al–Mg–Si and Al–Zn–Mg–Cu) were exposed to water-vapor-saturated air at 70°C. Further evidence is provided by the experimental observation for 99.95% purity aluminum artificial crevices exposed to high-purity water, where the local solution’s pH conditions within the crevices become acidic, following a similar time dependence as observed for most aluminum alloys exposed to aqueous saline solutions (87). On the basis of the above discussion, it is reasonable to conclude that the initial surface ‘‘oxides’’ chemically react with aqueous environments to form modified surface films that, in the absence of aggressive species, will exist in a state of equilibrium with the aqueous environment. Although not yet proven, it would seem reasonable that inhibitor anions 204 Holroyd such as chromate initially function by competitive adsorption on the initial oxide surface followed by the adsorbed species somehow modifying the subsequent film-formation process and the film properties. Augustynski’s (72) results from XPS experiments supports this view and he concluded that after the chromate anions adsorb on aluminum oxide surfaces, chemical reactions lead to hydrated chromium III and aluminum III oxides as well as chromium IV species being present in the protective film. Step 3: Effective Thinning of the ‘‘Oxide’’ Film by ‘‘Dissolution’’ or Anion Penetrating the Film The traditional view of the ‘‘oxide’’ film on aluminum participating in a corrosion reaction in an aqueous environment has been that it acts essentially as an ‘‘anhydrous inert barrier’’ (88,89). Despite all the evidence to the contrary, it is still common for the ‘‘oxide’’ to be considered as ‘‘chemically’’ inert in aqueous solutions with pH’s in the range 4–9. This is not an accurate view of the ‘‘oxide’’ films. A more realistic view of surface films participating in corrosion reactions should be of hydrated aluminum oxide layers, probably in a ‘‘colloidal state’’ (69) or as a semipermeable membrane (90). Unfortunately, these ideas of films existing in a colloidal state or as semipermeable membranes have not received the attention they deserve. An exception to this is the work conducted by Liepina and co-workers during the 1950s and 1960s (91–95), in which the corrosion process was envisaged in terms of colloidal–chemical effects occurring on the metal surface. For aluminum in an aqueous potassium chloride solution, these authors suggest the potential reaction sequence Al–AlCl3 –polyoxychloride intermediates– amorphous gels–boehmite–bayerite–hydrargillite, with the proviso that if the aluminum oxide is in a colloidal state, the chloride ion would peptize it and render it dispersible. [Interestingly, in 1967, Hoar (96) suggested that the strong negative charge caused by adsorbed anions could peptize oxide films. However, to the author’s knowledge, this idea was not pursued.] As will be discussed further in Step 4 of the proposed process, the modified ‘‘oxide’’ that is generated on aluminum surfaces exposed to aqueous environments must have a significantly higher ductility than those reported for anhydrous alumina films either stripped from or while adhered to an aluminum substrate (97). In the very early investigations of the pitting of aluminum in chloride solutions, it was believed that the chloride ion had the ability to penetrate the surface oxide (88,89), with the anion directly diffusing through the aluminum oxide lattice. This mechanism has now been discounted and Aluminum Alloys 205 the concept of ‘‘penetration’’ has been replaced with ones involving the following: 1. 2. Formation of soluble compounds or transitory species at critical sites (69,98,99) and\or The transport of species through a semipermeable membrane, formed as a result of the initial oxide chemically reacting with an aqueous environment (90) Step 4: Direct Attack of the Exposed Metal It is most improbable that the direct attack of ‘‘film-free’’ aluminum surfaces ever occurs during the corrosion of aluminum in aqueous environments: an exception perhaps being during stress corrosion (or corrosion fatigue) when the operative crack propagation mechanism, such as ‘‘hydrogen embrittlement,’’ involves periodic ‘‘brittle’’ crack jumps linking the main crack with subsurface internal cracks with surfaces that have not previously been exposed to the aqueous environment. Quantitative electrochemical data describing the behavior of ‘‘fresh’’ aluminum surfaces when exposed to aqueous environments is now available (100) based on data from a recently developed test method known as the guillotine method (101). Another exception is an air-stable phosphate surface directly bonded to pure aluminum without an intermediate oxide layer that has been generated by Rotole and Sherwood (102) under precisely controlled conditions generated during in situ experiments in ultrahigh-vacuum surface analytical equipment. Increasing experimental evidence now suggests that corrosion initiation and some corrosion processes can occur beneath the surface films that have formed as a result of the original surface ‘‘oxide’’ ‘‘chemically’’ reacting with its local environment (90,103). A good example of the above is the experimental evidence provided by Bargeron and Givens (104–107) and Bargeron and Benson (108), who exposed 99.999% aluminum with preanodized oxide surfaces to aqueous chloride solutions. In their initial studies, Bargeron and Givens (104) observed that small circular blisters (diameter ⬍ 100 µm) developed between the aluminum surface and the oxide when the systems electrochemical potential was given a sufficiently anodic short-duration pulse. Further work by Bargeron and Givens (105) for high-purity aluminum in 0.5M KCl indicated the following: 1. Blisters only formed if the steady-state electrochemical potential is above the pitting potential and the minimum potential for blisters dependence on the chloride concentration closely resembles that reported by others for the pitting potential (Fig. 20). 206 Holroyd Fig. 20 Blister growth potential given by Bargeron and Givens (105) compared with the pitting potentials for pure aluminum in various aqueous chloride solutions quoted by various authors. 䊐 from Ref. 120; 䉭 from Ref. 144. 2. Blister growth proceeds along the metal–oxide interface with the oxide exhibiting ‘‘plastic’’ behavior. 3. Blisters generally are circular, but the shape can be locally distorted by grain boundaries and other microstructural features. 4. Corrosion initiates locally in a peripheral zone toward the blister’s perimeters, prior to the blister bursting due to excessive internal gas pressure. 5. An ‘‘oxide’’ skin of some description remains in place over the ‘‘ruptured’’ blister which is capable of maintaining an ‘‘occluded cell’’ and a localized ‘‘solution chemistry.’’ 6. Gas compositions generated from the blister/pits are specific to the anions in the aqueous environment; for instance, Bargeron and Benson (108) detected hydrogen exclusively for chloride solutions and de- Aluminum Alloys 207 tected a range of gases including hydrogen, nitrogen, and nitrous oxide for nitrate solutions. Further work by Bargeron and Givens (106) matched the hydrogen-bubblegeneration activity to the current fluctuations recorded during potentiostatic experiments on aluminum specimens with a single active pit immersed in an aqueous 1M KCl solution. They claimed that irrespective of bubble rate or size, the current noise was directly related to bubble generation. To date, attempts to generate and observe blister formation on pure aluminum surfaces with air-formed oxide films have failed unless the films ˚ . Despite this, Bargeron and are artificially thickened to around 100 A Givens (105) and others (109) firmly believe that the blistering process is the precursor to the pitting of aluminum alloys. Sound evidence in favor of this argument include the following: 1. 2. Current transients similar to those found when blisters are observed are also detected during pit initiation when blisters are not resolvable. Natishan and McCafferty (109) have observed blister formation in ˚ -thick oxide layers on ion-implanted pure aluminum surfaces 30-A containing 4 or 12 at% Cr, Si, Zr, Nb, Mo, or Cr ⫹ Mo. (In these cases, it was suggested that ion implantation somehow modifies the properties of the ‘‘oxide’’ forming after the reaction with the aqueous environment, thereby allowing larger blisters which are detectable to form.) 1. Influence of Electrochemical Potential on Corrosion Initiation The conventionally accepted view, as proposed by Pourbaix and his co-workers (53,110), is that two characteristic potentials are associated with the pitting corrosion of aluminum; namely the critical (or breakdown) potential Ep (a potential below which pits cannot activate) and the protection (or repassivation) potential ER; a more active potential below which pre-existing pits will no longer remain active or grow); see Fig. 21. (Throughout this chapter, these two potentials will be referred to as Ep, the pitting potential, and ER, the repassivation potential.) The accepted view was challenged in the early 1970s when results from studies conducted at the University of Trondheim in Norway (111–113) suggested that the potentials Ep and ER were identical and, therefore, there was only one characteristic potential associated with the pitting corrosion of aluminum in saline solutions. In these detailed studies, Broli and Holtan (111) measured Ep 208 Holroyd Fig. 21 Schematic of a typical anodic polarization curve for aluminum and its alloys in a saline environment. and ER values for 1S aluminum (99.0–99.3% purity aluminum) in de-aerated 3% NaCl using three electrochemical methods: (a) potentiodynamic (1–100 mV/ min), (b) a quasi-stationary method, and (c) a potentiostatic method typically lasting 4–6 days [later adding galvano-kinetic polarization as a fourth method (113)]. During further detailed studies at the University of Trondheim by Nisancioglu and Holtan (114,115), it was realized that some of the results published by Broli and Holtan were ‘‘flawed’’ due to crevice conditions being established during some of the electrochemical experiments. Once the experimental problems were resolved, two definitive papers emerged in 1978 by Nisancioglu and Holtan (114,115) and conclusions from these studies for 1S aluminum (99.0–99.3% purity aluminum) in de-aerated 3% NaCl at 30°C may be summarized as follows: The critical pitting potential, Ep (114): 1. It is a fairly reproducible quantity for aluminum in saline solutions if measured using ‘‘static’’ electrochemical methods. However, high- Aluminum Alloys 2. 3. 209 potential scan rates during potentiodynamic polarization methods can lead to appreciable errors, and scan rates less than a few millivolts per minute are recommended. It is a strong function of the aqueous solution’s chloride concentration and is independent of surface treatment, film thickness, and solution stirring rates up to 2040 rpm. When measured by a potentiostatic method and is typically ⫺0.76 V (SCE), with • • • • Pits Pits Pits Pits almost never initiating at E ⬍ ⫺0.76 V (SCE) seldom initiating at E ⫽ ⫺0.76 V (SCE) often initiating at E ⫽ ⫺0.75 V (SCE) always initiating at E ⬎ ⫺0.75 V (SCE) 4. It is not an ideal criterion to determine an aluminum alloy’s pitting susceptibility. The protection potential, ER (115): 1. It is typically about 100 mV more active than Ep. 2. It is the lowest potential for pit growth. Pits grown at potentials above ER but below Ep involve growth from pre-existing occluded microscopic sites via crystallographic pitting or tunneling, whereas pitting at potentials above Ep is generally macroscopic in nature. 3. It may be explained in terms of active chloride participation in competition with the hydroxide in the metal dissolution reaction (116), whereas the basic factors determining the critical pitting potential, Ep, remain unclear, although its measurement is reasonably reproducible under a given set of experimental conditions. 2. Ep Dependence on ‘‘Dynamic’’ Versus ‘‘Static’’ Polarization: ‘‘Induction’’ and ‘‘Statistical’’ Effects Differences between Ep measurements from ‘‘dynamic’’ and ‘‘static’’ polarization methods are mainly associated with the ‘‘induction’’ time needed for pits to initiate. Several authors have studied this phenomenon using various electrochemical techniques, and reviews are provided elsewhere (117,118). Foroulis and Thubrikar in the mid-1970s (98,119,120) evaluated the influence of various factors on Ep values for 99.99% pure aluminum in de-aerated aqueous potassium chloride solutions with concentrations in the range 0.01M– 3M and temperatures in the range 5–60°C. In these studies, the aluminum surfaces were initially mechanically polished to produce a mirror finish and then anodized using one of two methods to produce oxide films of known thickness. Critical pitting potentials were determined using two methods, one potentiostatic and the other a quasi-potentiodynamic method. In a typical potentiostatic 210 Holroyd experiment, the electrode potential would be increased anodically by a potential step of 10–20 mV and then held constant with the current continuously monitored for up to 24 h if a significant current increase was not recorded. After this, the aluminum electrode surface was examined for pit initiation using a low-power optical microscope. The critical pitting potential, Ep, from these experiments was defined as the lowest potential that the pits were optically observed, which turned out also to be the lowest potential where the current significantly increased after the potential increment. A typical result from a ‘‘static’’ polarization experiment is given in Fig. 22, which shows that for the particular conditions used, Ep is 0.81 V (SCE) and that the current transient at Ep has an ‘‘induction’’ time, τ, of about 8 min during which the current remains negligible prior to it increasing significantly at longer times (98,119,120). ‘‘Induction’’-time effects lead to Ep values from potentiodynamic polariza- Fig. 22 Typical potentiostatic polarization curves for pre-anodized 99.99% pure aluminum in 3M KCl. [Note the 8-min induction time for pitting at E ⫽ ⫺0.81 V (SCE).] (Data from Ref. 120.) Aluminum Alloys 211 tion studies being more noble than values obtained from potentiostatic methods. The reason for this is that by the time the induction time has elapsed, the electrochemical potential in a ‘‘dynamic’’ experiment has moved on from the ‘‘potentiostatic’’ Ep to a more positive value. Typical examples of this effect are provided by the Ep data published by Foroulis and Thubrikar (119,120) from potentiostatic and potentiodynamic polarization experiments on pure aluminum (99.99%) in aqueous potassium chloride solutions at various chloride concentrations for a given temperature and at various temperatures for a given solution chloride concentration. These data are reproduced in Figs. 23 and 24, along with comparable data from other studies. ‘‘Induction’’ times are a common feature of pit initiation under potentiostatic test conditions. Published data indicate that it is a statistical variable (121,122) with ‘‘induction’’ times, τ: • Decreasing with increasing chloride concentration (98,122), potential (112,122), or temperature (98,122), Fig. 23 Pitting potentials for pure aluminum as a function of chloride concentration given by potentiostatic and potentiodynamic polarization methods. 䊊, 䊉 from Ref. 120; ■ from Ref. 123; ⫻ from Ref. 127; 䉮 from Ref. 126; 䉭 from Ref. 125. 212 Holroyd Fig. 24 Effect of temperature on the pitting potential for pure aluminum in a 3M KCl solution given by potentiostatic and potentiodynamic polarization methods. 䊊, 䊉 from Ref. 120; 䉭 from Ref. 123. • Being independent of solution pH (98) or chloride concentrations below ⬃0.5M (98) ˚ (98) or being to• Increasing with oxide film thickness above ⬃1300 A tally independent (123) Shibata and Sudo’s study on pure aluminum (99.99%) in de-aerated 3.5% NaCl (121), in addition to giving further data on potential sweep rate effects on Ep measurements from anodic potentiodynamic polarization experiments, also provides excellent examples of the statistical nature of Ep measurements. Measured Ep values in the form of cumulative probability plots for various potential sweep rates and mean Ep values as a function of potential sweep rate for various temperatures (298–323 K) are shown in Figs. 25 and 26. From these data, it is apparent that the absolute value and the statistical spread of the Ep measurements both increase with increasing potential scan rate, and the effects become less dominant as the solution temperature increases. Similar statistical data have recently been published by Sato and Newman (124) for pure aluminum (99.999%) in a 0.5M NaCl solution at room temperature. Aluminum Alloys 213 Fig. 25 Experimental pitting potentials determined obtained using various anodic potential sweep rates for pure aluminum in de-aerated 3.5% NaCl and plotted on normal probability paper. (Data from Ref. 121.) 3. Influence of Chloride Concentration, Solution pH, and Temperature and the Nature and Thickness of Oxides on Ep and ER a. Chloride Concentration Although, it is true that ‘‘static’’ polarization methods provide more accurate Ep and ER measurements than ‘‘dynamic’’ methods, for convenience most researches have employed ‘‘dynamic’’ polarization methods. No significant problems arise as long as ‘‘dynamic data’’ are used to evaluate trends rather than making absolute judgments. With this in mind, we can use published polarization data to characterize the effect of a range of variables on the characteristic pitting potentials for pure aluminum in aqueous environments. The dependence of Ep on chloride concentration (shown in Fig. 23) is well known (120,123,125–127) and several authors have published relationships, for instance (120): Ep ⫽ ⫺0.80 ⫹ 0.120 log[Cl⫺] V (SCE) 214 Holroyd Fig. 26 Median values of the pitting potentials for pure aluminum in a de-aerated 3.5% NaCl solution at various temperatures, plotted as a function of the square root of the anodic polarization potential sweep rate. (Data from Ref. 121.) b. Solution pH Until recently, it was generally accepted that Ep and ER values are effectively insensitive to the solution’s pH for pH’s in the range 4– 9. More recently, it has been recognized that the local solution’s pH changes can occur at aluminum surfaces during the electrochemical experiments used to measure Ep and ER values and when these changes are compensated for the resultant Ep and ER values can be influenced (128). Lampeas and Koutsoukos (128) showed that the cathodic polarization characteristics of 99.99% purity aluminum in aqueous saline solutions are relatively insensitive to solution pH changes, whereas those for anodic polarization are sensitive in the pH range 6–7, where films are more protective. Interpretation of these data with respect to the initiation of pitting needs further evaluation, particularly if initiation occurs beneath hydrated surface films, as has been suggested by several researchers (90,103–109). c. Temperature Potentiodynamic anodic polarization experiments on 99% purity aluminum (AA1100) and a range of other aluminum alloys were conducted in the de-aerated synthetic seawater at various temperatures in the range 25–150°C during the early 1970s in a program to evaluate the suitability of the alloys for use in desalination applications (129,130). Findings from these studies were as follows: 1. Free-corrosion potentials increase over the first 40 h or so exposure to Aluminum Alloys 215 the aqueous environments with the stabilizing values increasing with increasing temperature 2. Repassivation potentials, ER, are insensitive to soak time and decrease linearly with increasing temperature, from ⫺760 mV at 25°C to ⫺800 mV (SCE) at 150°C 3. Critical pitting potentials, Ep, increase with soak time and the stabilization times and potential shifts increase significantly at higher solution temperatures. The observed temperature dependencies for free-corrosion potentials (FCP), Ep and ER, are understandable if the latter potential is deemed independent of surface filming characteristics, as opposed to the other two parameters being considered highly dependent on surface filming characteristics. This is justifiable because (a) the FCP is associated with establishing and maintaining a surface film, (b) Ep is associated with the local breakdown of surface films, and (c) ER is defined as the potential below which activity no longer occurs when the local film has previously already suffered breakdown. B. Aluminum Alloys: Oxide Development and Corrosion Initiation in Aqueous Environments Effects alloying additions have on the filming characteristics of aluminum alloys in aqueous environments and any subsequent corrosion behavior are dependent on the additions themselves, the actual microstructural form(s) adopted in a given alloy-temper condition, and the prevailing local environmental conditions for film formation. Considerable research effort has been directed toward answering the question, How do alloying additions modify the inherent corrosion resistance of aluminum alloys? Suggested mechanisms include the following: • The intermetallic providing local anodic and/or cathodic sites • Generation of oxide ‘‘weak spots’’ and oxide flaws associated with second phases • Alloying additions entering and modifying the properties of passive layers • Modified local solution chemistries developing within corrosion sites influencing growth kinetics The extremes a given alloying addition may introduce into an alloy microstructure are that the addition may be effectively insoluble in aluminum and therefore be present only as a second phase or an addition may be highly soluble and be present in a single-phase solid solution. The most common situation for most major alloying additions in commercial aluminum alloys (e.g., Cu, Mg, Mn, and Li) is that the additions have a limited solid solubility in aluminum, and at low 216 Holroyd concentrations, they are present in the form of a solid solution, whereas at higher concentrations, they are also present as a second phase. For heat-treatable aluminum alloys, the situation is further complicated by the fact that heat treatment may modify the microstructural form adopted in the alloy, and for non-heattreatable alloys, the same may be promoted by cold or warm working processes. In view of the above situation, it is not too surprising that the precipitation of electrochemically distinguishable second phases (e.g., CuAl2 in Al–Cu alloys and Mg3Al2 in Al–Mg alloys) at preferential microstructural sites means that alloys can suffer additional forms of localized corrosion to the pitting suffered by unalloyed aluminum. These various forms of localized corrosion will be covered in the following discussion below, which will initially focus on binary aluminum alloys, move on to ternary alloys, and, finally, deal with the more complicated commercial aluminum alloys that contain multiple deliberate alloying additions along with others present as impurities. C. Binary Alloys Many studies have been published on the corrosion of binary aluminum alloys in aqueous environments. Initial studies concentrated on binary systems involving the elements used as major alloying additions (i.e., Al–Cu, Al–Mg, Al–Zn, Al– Mn, and Al–Li). In more recent times, in response to the goal of generating a ‘‘stainless’’ aluminum alloy, corrosion studies have involved binary aluminum alloys bases containing elements known to enhance the passivity of stainless steels (e.g., Cr, Mo, Ta, Zr, Nb, and W). Because these elements all exhibit extremely low solid solubilities in aluminum (typically less than 1 at%), these researchers have had to employ novel techniques such as rapid solidification processes (e.g., splutter deposition or melt spinning) or ion implantation to generate the alloys. Detailed studies have also been conducted on various binary alloys containing alloying elements that activate aluminum (e.g., Hg, In, Sn, Zn) (57,131– 133). This aspect of aluminum alloy corrosion will not be discussed in detail here, as our focus is on aluminum alloys used in structural and engineering applications rather than those employed as sacrificial anodes or battery anodes. The earliest corrosion studies on aluminum binary alloys not surprisingly were conducted on the Al–Cu system, which, after being discovered in Germany in 1906, resulted in the alloy duralumin and the first commercial exploitation of an age-hardenable aluminum alloy in 1908 (134). It has been appreciated since the early 1940s that the copper in a Al–4% Cu (wt%) binary alloy may be retained in solid solution if the alloy is solution heat treated and sufficiently rapidly quenched. In this metallurgical state, the dissolved copper is thermodynamically unstable and will preferentially precipitate as CuAl2 at grain boundaries if the alloy is exposed to temperatures above around 120°C or will have already precipi- Aluminum Alloys 217 tated to some extent if the quench rate employed after solution heat treatment was insufficient. This miscrostructural condition promotes a susceptibility to intergranular corrosion and researchers quickly realized that the localized corrosion was related to the grain-boundary precipitation of CuAl2 (135,136). Dix and coworkers developed an electrochemical based model for the intergranular corrosion promoted in these situations. In this model (137), they suggest that the intergranular corrosion occurs along a narrow zone adjacent to the grain boundary that is relatively depleted in copper and anodic to both the main body of the grain and the CuAl2 precipitates on the grain boundary. (A schematic of the local microstructure is given in Fig. 27.) The weakness of the Dix et al. model for intergranular corrosion as it was initially presented is that it relies solely upon whether a given phase is anodic or cathodic with respect to another phase when exposed to a saline solution doped with hydrogen peroxide. No consideration is given to the magnitude of the current that would flow in a local cell or whether the poly-phase system’s free-corrosion potential would be greater or smaller than that needed to promote local corrosion. Ideas presented by Galvele and his co-workers (138–140) show that the Dix et al. model for the intergranular corrosion of Al–Cu alloys can be justified if it includes a consideration of the breakdown potentials of the phases present in the system (141). Fig. 27 Schematic representation of an aged Al–Cu (wt%) alloy microstructure. 218 Holroyd An excellent overview of the corrosion behavior of Al–Cu binary alloys in aqueous saline environments is provided by Muller and Galvele’s studies (142). They established that a clear relationship existed between alloy microstructure, electrode potential, and corrosion behavior in de-aerated 1M sodium chloride at 25°C as a function of alloy composition and thermal aging. The effect of copper content on the pitting potentials of Al–Cu solid solutions is shown in Fig. 28 and the effect that aging time at 240°C has on the pitting potential, alloy hardness, and corrosion behavior for an Al–3.3% Cu (wt%) alloy as a function of electrode potential and the calculated copper content of the solute depleted zone associated with the grain boundary is shown in Fig. 29. From these data, Fig. 28 Pitting potentials for various aluminum binary alloys in a de-aerated 1M NaCl solution at 25°C. ■ from Ref. 140; 䉱 from Ref. 140; 䊉 from Ref. 142. Aluminum Alloys 219 Fig. 29 Corrosion mode, pitting potential, and alloy hardness for an A1–3.3% Cu (wt%) as a function of aging time at 240°C. (Data from Ref. 142.) it should be evident that the observed corrosion mode is directly related to an alloy’s microstructure and the prevailing electrochemical potential. For instance: 1. For an as-solution treated and rapidly quenched alloy microstructure, the alloy is passive at potentials below the pitting potential and pitting only occurs at higher potentials. 2. For aging times where the pitting potential for the grain-boundary regions is below that for the grain interiors, a potential regime will exist where only intergranular corrosion occurs (Fig. 29). 3. After long aging times, the potential domain for intergranular corrosion disappears (Fig. 29). 220 Holroyd Based on the above, it would seem reasonable that suitably heat-treated Al–Cu binary alloys should display two pitting potentials during anodic polarization studies: one for the grain-boundary region and a second at a higher electrode potential for the grain interiors. Such results have been reported by Urushino and Sugimoto (143) for an Al–4 Cu alloy aged at 170°C and polarized in a de-aerated 1M NaCl solution with its pH adjusted to 10. These results are reproduced in Fig. 30 and the good fit with the Muller and Galvele data is obvious. Muller and Galvele’s corrosion studies also included work on the Al–Mg and Al–Zn binary-alloy solid solutions (140). Significant differences were readily apparent between the various alloy systems. Alloying additions up to around 5% (wt) led to pitting potentials in an aqueous 1M NaCl solution displaying positive shifts for Al–Cu alloys, small effects for Al–Mg alloys, and shifts in the negative direction for Al–Zn alloys (Fig. 28). A further significant difference noted was the electrochemical potentials for the intermetallic phases precipitating from supersaturated solid solutions and those for the denuded zones adjacent to these precipitates. In the case of Al–Cu alloys, the intergranular precipitation of Al2Cu- Fig. 30 Effect of aging time at 170°C on the grain boundary and matrix pitting potentials for an A1–4 Cu (wt%) in a deaerated 1M NaCl solution with its pH adjusted to 10. (Data from Ref. 143.) Aluminum Alloys 221 generated adjacent regions that were relatively anodic, whereas for the Mg3Al2 precipitation in Al–Mg alloys, the precipitate itself was relatively anodic; see data given in Table 2. Pitting attack morphologies for the Al–Mg alloys were crystallographic, closely resembling that previously reported by Galvele and de Micheli (144) for pure aluminum, and there was no evidence of any magnesium buildup within the pits. The pitting attack for the Al–Zn alloys was dependent on an alloy’s zinc concentration, and microprobe studies indicated that zinc accumulation occurred within pits. For low zinc levels, the pitting was crystallographic; however, as the zinc level increased, the pits became more irregular and more tunnellike, probably with some subsurface propagation (140). Based on results from their experiments involving scratching oxide surfaces during anodic polarization experiments, Muller and Galvele (140) concluded that the oxides forming on various aluminum alloys had different properties. The presence of oxide films on pure aluminum, Al–Cu, and Al–Zn did not interfere with or accelerate the initiation of pitting, whereas films forming on the Al–Mg alloys acted as a barrier to the pitting process, with pits for these alloys, unlike the others, nucleating preferentially on the scratch lines. Muller and Galvele (140) noted during their polarization studies on the higher magnesium binary alloys (Al–3 Mg and Al–5 Mg) that the measured currents became unsteady at potentials just below the pitting potentials. The ob- Table 2 Published Pitting Potentials, Ep, for Various Aluminum Binary Alloys and Aluminum-Based Precipitates in De-aerated 1M NaCl at 25°C and 53 g/L NaCl ⫹ 3 g/L H2O2 Alloy Pitting potential of de-aerated 1M NaCl [V (SCE)] (Ref.) Al–2 Cu Al–4 Cu ⫺0.774 ⫺0.777 ⫺0.654 ⫺0.594 Al–1 Zn Al–3 Zn Al–3 Mg Al–5 Mg Al2Cu Al3Mg2 ⫺0.864 (140) ⫺0.994 (140) ⫺0.814 (140) ⫺0.824 (140) ⫺0.654 (144) ⫺0.994a Pure aluminum a Extrapolated from data given in Ref. 145. (144) (145) (142) (142) Pitting potential 53 g/L NaCl ⫹ 3 g/L H2O2 [V (SCE)] (Ref.) ⫺0.764 (144) ⫺0.644 (141) ⫺0.598 (135) ⫺0.604 (141) ⫺0.874 (141) ⫺0.994 (141) ⫺0.784 (141) ⫺0.794 (141) ⫺0.644 (6, 144) ⫺0.978 (135) 222 Holroyd served transient peak currents were two to three times those needed to maintain the passive film and occurred at increasing frequencies as the potential approached the alloy’s pitting potential. These authors did not comment on whether these effects also occurred during their experiments on Al–Zn binary alloys; however, as will be discussed later, these effects are now known to occur for Al– Zn binary alloys in aqueous saline solutions (124,145). Perhaps the most significant outcome from the above studies by Galvele and his co-workers is that the results are consistent with his previously proposal pitting model, where the pitting potential is the minimum potential for sufficient localized acidification to be maintained at the metal–solution interface within a pit (138). According to Galvele (138), the pitting potential is given by Ep ⫽ E p* ⫹ η ⫹ ϕ ⫹ Einh (1) where Ep is the pitting potential, E p* is the pitting potential of the alloy in a pitlike environment, η is the overpotential necessary to draw a net anodic current through the pit, ϕ is the potential gradient through the pit, and Einh is the contribution due to inhibitors present in the solution. Substitution of appropriate values for 99.99% purity aluminum, Al–3 Cu, Al–3 Mg, and Al–3 Zn into Eq. (1) yields the results given in Table 3. The predicted pitting potentials are in excellent agreement with those determined experimentally. [Values for E p* are based the alloy’s corrosion potential measured in a saturated AlCl3 solution after 2 h exposure (140). Values for η were taken as the polarization above the corrosion potential necessary to maintain an alloy’s anodic current density at 1 mA/cm2 in a saturated AlCl3 solution; see Table 3. ϕ for a 1M NaCl solution is approximately 0.050 V (138) and because there are no inhibitors Einh ⫽ 0.] Sato and Newman (145) have recently published elegant experimental studies on the role zinc additions up to 5% (wt) play during the activation of pitting corrosion for high-purity binary Al–Zn alloys in a de-aerated aqueous 0.5M NaCl solution. Their results are consistent with Zn effects on the pitting potential being the following: Table 3 Comparison of Pitting Potentials Calculated Using Eq. (1) with Experimental Values Obtained in De-aerated 1M NaCl at 25°C Alloy 99.99% aluminum Al–3 Cu Al–3 Mg Al–3 Zn E *p η φ ETheory Ep ⫺1.00 ⫺0.76 ⫺0.10 ⫺1.06 ⫹0.17 ⫹0.04 ⫹0.15 ⫹0.01 ⫹0.05 ⫹0.05 ⫹0.05 ⫹0.05 ⫺0.78 ⫺0.67 ⫺0.80 ⫺1.0 ⫺0.77 ⫺0.64 ⫺0.81 to ⫺0.77 ⫺0.99 Source: Data from Refs. 140 and 142. Aluminum Alloys 223 1. Predictable from Zn’s effects on dissolution kinetics in pitlike environments 2. A ‘‘pure’’ Galvele effect adhering to his local acidification model of the pitting potential 3. Independent of zinc needing to modifying properties of the surface films Furthermore, their results clearly showed the following: 1. The same number of potential pit sites exist irrespective of a binary alloy’s zinc level (presumably they depend on the nature of the oxide and sites generated with time by passive dissolution). 2. Zinc additions have little or no effect on pit nucleation or its frequency, but effects are solely associated with pit propagation by the enhancement of local dissolution kinetics within the active pits. 3. Zinc has no direct effect on the local film breakdown and repair events occurring at potentials either just below or just above the measured pitting potential. 4. Zinc’s role is to act as an activator during the propagation stage of pitting making pits grow faster and, in the case of metastable pitting, allowing pits to exist for longer times and in some circumstances grow sufficiently to aid their transition to stability. Justification of the above statements for the Al–Zn binary alloys containing up to 0.13% Zn (wt) was that the pitting potentials and repassivation potentials for the binary alloys were statistically indistinguishable from those for 99.999% purity aluminum (Fig. 31), as was the frequency of the metastable pitting events as a function of exposure time in the saline solution. (The 0.13% Zn addition, however, did extend the lifetime of the pits and so zinc was activating aluminum dissolution.) Justification for the higher-zinc-containing binary alloys was not possible by the direct comparison of pit nucleation frequency with those for pure aluminum at a given potential because no events occur on pure aluminum at the lower potentials associated with the pitting potentials of the higher-zinc-containing binary alloys. However, pit nucleation behavior was compared with pure aluminum by measuring the standard deviation of the passive current, which indicated passive film breakdown associated with pit nucleation in the presence of the chloride ion (145). In these experiments, the authors showed that the current ‘‘noise’’ for pure aluminum exposed to an aqueous sodium borate solution containing 0.5M NaCl decreased to that of the borate solution without the chloride addition when the potential was decreased below the pitting potential of a Al– 2 Zn alloy. This, the authors claim, successfully confirms that pitting of Al–Zn alloys is the result of nucleation events that occur with about the same potential dependence and possibly the same frequency as on pure aluminum. 224 Holroyd Fig. 31 Normal probability plot of the pitting potential for superpure (99.999%) aluminum and A1–0.13 Zn (wt%) in a de-aerated aqueous 0.5M NaCl solution. (Data from Ref. 124.) Aluminum Alloys 225 Corrosion studies on Al–Li-based alloys increased when the aerospace industry became interested in the alloy system’s potential promise of a 10% increases in stiffness combined with a 10% density reduction over the conventional high-strength aluminum alloys (146). Published work on Al–Li binary alloys indicates that the corrosion resistance in saline solutions is no worse than that observed for pure aluminum. Corrosion occurs as crystallographic pitting and is insensitive to thermal aging, irrespective of whether the Al–Li binary alloy is single phase, with the lithium totally in solid solution, or the lithium also present in the alloy as the normal age-hardening precipitate, δ′, Al3Li (147,148). The only exception is when Al–Li binary alloys are subjected to extended overaging heat treatments and enhanced corrosion rates are observed with pits initiating and spreading from grain boundaries (147). The latter has been attributed to the extended overaging promoting grain-boundary precipitation of δ, Al–Li, a highly anodic phase (149,150). The reason why, unlike Al–Cu and Al–Mg binary alloys, the corrosion behavior of Al–Li binary alloys are insensitive to the alloy addition being either in solid solution or precipitated as a hardening phase is twofold. First, the Al3Li phase promoting age hardening is a coherent phase and the electrochemical potential differences between it, the adjacent solute-depleted zones, and the grain interiors are believed to be minimal and, second, lithium dissolution can lead to beneficial modifications to the surface films generated in aqueous environments. Over the last 20 years, attempts have been made to develop ‘‘stainless’’ aluminum alloys that will offer high resistances to localized corrosion in chloride environments. The approach adopted has been to use nonequilibrium methods that will simultaneously do the following: 1. Eliminate the second phases, usually present in conventionally cast alloys, that are known to be detrimental to corrosion resistance 2. Produce supersaturated alloys containing elements thought to enhance passivity Nonequilibrium methods employed to introduce a wide range of elements in aluminum include ion implantation (109,151–153) and rapid solidification methods such as melt-spinning (154) and sputter deposition (155–161). One of the earliest studies involved using ion implantation to introduce Mo⫹ ions or Ar⫹ ions into 99.99% pure aluminum and a high-strength Al–Zn– Mg–Cu alloy, 7075-T6 (151). Implantation of Mo⫹ ions into the pure aluminum promoted a measurable improvement in the pitting corrosion resistance in an aqueous saline environment with the pitting potential, Ep, increasing by around 200 mV, whereas for 7075-T6, the effect was less pronounced with the Ep only increasing by around 100 mV. The authors suggested that these beneficial effects were due to either the incorporation of Mo into the passive film or the reprecipitation of some Mo-containing species on the passive film. Effects due to Ar⫹ ion 226 Holroyd implantation into either pure aluminum or the 7075-T6 were limited to small increases in the measured Ep values which were attributed to a slight thickening of the air-formed oxide films induced by the ion implantation. Corrosion data in de-aerated 0.5N NaCl has been provided by Yoshioka et al. (154) for a wide matrix of rapidly solidified aluminum binary alloys containing 2, 4, and 6 (at%) Mg, Ti, Mn, Cr, Fe, Ni, Cu, Zn, Zr, and Si produced by a meltspinning process. Pitting potentials for all the alloys were significantly ennobled (and for pure aluminum) except for the alloys containing Mg, Fe, and Zn. Corrosion rates at the free-corrosion potential showed similar trends that reflected the observed decreases in both the anodic and cathodic current densities. These beneficial effects were attributed to both the supersaturation of the aluminum solidsolution phase with solute atoms and to the elimination or decrease in the number of intermetallic phases available to initiate pitting. Extensive work studying the evolution of passive film chemistry on a range of sputter-deposited Al–X binary alloys, where X included Cr, Ta, Mo, and W, was conducted at the Martin Marietta Research Laboratories by Moshier and coworkers during the mid-1980s and the early 1990s (155–159). Beneficial effects on pitting resistance promoted by the various additions should be obvious from Fig. 32, showing a compilation of anodic polarization curves from these studies for pure aluminum and a range of the supersaturated aluminum binary alloys obtained in 0.1M NaCl. X-ray photospectroscopy (XPS) studies on the passive films formed on Al–Mo, Al–Cr, and Al–Ta alloys after exposure to a 0.1M NaCl solution revealed that the films contained 5–10% of the oxidizing solute as MoO4⫺, CrOOH, and Ta2O5, respectively, with the concentrations increasing fivefold as the overpotential increased from the open-circuit potential to the pitting potential (156,157). On the basis of these results, the authors proposed that the enhanced pitting corrosion resistance resulted from the above species somehow impeding chloride ingress through the passive film. A different explanation was invoked for Al–W alloys because although W was very effective in improving the pitting resistance (Fig. 32), the XPS studies failed to detected any W in the passive films. Here, it was proposed that an undefined synergistic effect between W and Al2O3 somehow led to a more stable oxide layer at the metal–oxide interface (158). Natisham et al. (152) offer an alternative hypothesis to account for the beneficial effects on the pitting corrosion resistance found for a range of aluminum binary alloys produced by an ion-implantation method. Here, the authors assume that the rate-determining step (RDS) for pitting is the adsorption of chloride ions on to an alloy’s filmed surface and, hence, it is critical that the oxide film covering the alloy must be positively charged so that chloride ions will be adsorbed. As evidence supporting this proposal, they presented a correlation between the pitting potentials for Ta, Ti, Cr, Cu, Fe, Al, Zn, and Mg and the pHZCH Aluminum Alloys 227 Fig. 32 Anodic polarization curves for various sputter-deposited aluminum binary alloys in an aerated aqueous 0.1M NaCl solution. (Data from Ref. 158.) (pH of zero charge) for the metals hydrated oxide in an aqueous 1N NaCl solution (152). More recent studies have questioned the validity of Natisham et al.’s mechanistic interpretation and the relevance of the correlation between the pitting potential and pHZCH. Perhaps the most convincing of these arguments is provided by the work of Vijh (162), who suggests that the correlation between Ep and pHZCH exists because a linear relationship also exists between pHZCH and the metal–metal bond energy b(M–M) (163). Vijh suggests that the critical relationship is the one between Ep and b(M–M) (Fig. 33) and not the one cited by Natisham et al. The acceptance of b(M–M) as the dominant parameter rather than pHZCH is relatively straightforward if the RDS for pitting is attributed to a process such as the devel- 228 Holroyd Fig. 33 A plot of pitting potentials of various aluminum alloys against the metal–metal bond energy b(M–M) values of the alloying element in the binary alloy. (Data from Ref. 162.) opment of appropriate local solution conditions within a pit rather than one involving an aggressive species’ surface adsorption on to or its transport through a passive film. Szklarska-Smialowska (164) has proposed yet another mechanistic interpretation. The RDS for stable pit growth in her model is deemed to be the establishment of an appropriate solution chemistry within a pre-existing defect in the film that can support active dissolution of the metal at the metal–oxide interface. This is a variant of the Galvele model for pitting (138), with the main difference being an extra caveat in that the envisioned role for the alloying addition is to establish the level of acidification needed to dissolve the alloying addition’s oxide Aluminum Alloys 229 in the pit environment. For instance, in this model, an alloying addition can either cause Ep to increase by providing an oxide that is less soluble than alumina in an acidified pit environment (e.g., Cr and W oxides) or to decrease Ep by providing an oxide that is more soluble. Justification for the passive films themselves not being involved in the RDS is based on the following premise: 1. The observation that the passive current densities for pure aluminum and Al–X alloys in a given saline environment are effectively identical (see data in Fig. 32) and, hence, so are the physical and/or electrical properties of the passive films. 2. The passive films forming on pure aluminum in aqueous saline environments are not effective obstacles to the penetration of chloride ions and water to the metal surface (165) 3. Support for item 2 is provided by the immediate generation of electrochemical current noise when aluminum is immersed in chloride solutions as reported by Uruchurtu and Dawson (166) Inturi and Szklarska-Smialowska (165) used the model to explain the beneficial effects on the pitting characteristics of sputter-deposited Al–4 Cr, Al–9 Cr, Al– 10 Ta, and Al–17 Ta (at%) alloys exposed to a de-aerated 0.1M NaCl solution. The observed open-circuit potentials (OCP) and the passive current densities for these alloys were similar to pure aluminum’s, other than the OCP for the Al– 17 Ta alloy, which was more positive. This was attributed to the Al–17 Ta alloy’s air-formed oxide containing Ta5⫹ species as well as Al3⫹ species, whereas Auger electron spectroscopy only detected Al3⫹ cationic species in the air-formed oxide films on the other alloys. Vijh (162) has suggested that the correlative trend between the aluminum binary-alloy pitting potentials and the solubility of the alloying element oxide quoted by Szklarska-Smialowska (164) is heavily weighted by the presence of W and Ta among the nine elements studies. He presents an alternative interpretation of the data presented by Inturi and Szklarska-Smialowska (165) based on a Galvele-type model with active dissolution of bare metal initiating in Richardson and Wood (81,82) type of microfissures and defects in the oxide films. In support of his interpretation, Vijh demonstrates that the pitting potentials of the binary aluminum alloys quoted by Inturi and Szklarska-Smialowska (165) show the expected relationship with the solid-state cohesion metal-to-metal bond energies b(M–M) of the alloying element (163); see Fig. 33. Some of the controversies outlined above may be reconciled using data and ideas from a study on the pitting resistance of sputter-deposited aluminum alloys presented by Frankel et al. (161). These authors used thin-film sputter˚ thick) and measured pitting potentials, Ep, and deposited alloys (1000–2000 A the repassivation potentials, ER, in an aerated aqueous 0.1M NaCl solution. The 230 Holroyd alloys evaluated included Al–Nb, Al–Mo, Al–Cr, and Al–W with solute concentrations up to around 35 (at%). A major advantage for the Frankel et al. study was that the thin films used in the corrosion studies led to two-dimensional pits that accurately simulate the behavior of small three-dimensional pits (167). This overcomes the usual problems encountered with pit depth and ohmic potential drop and/or diffusional path issues and facilitates the generation of highly reproducible repassivation potentials, ER (167). Important findings from this study include the following: 1. For a given alloy system, the anodic pit current densities just prior to repassivation do not increase with the pitting resistance or an alloy’s solute content. 2. Although the pitting potentials of many of the sputter-deposited alloys were significantly higher than that for pure aluminum, the Ep for a given alloy was only slightly higher than its repassivation potential. Based on these findings, the authors concluded that although alloy solute enrichment in the alloys influenced the local environmental conditions needed for pit growth, any effects on the passive films themselves were relatively minor, even when the beneficial effects due to film modification from long-term ‘‘aging’’ (laboratory air exposure for 4 years), displayed by some of the alloying systems, was taken into account. In essence, the work by Frankel et al. provides a link between the nonequilibrium aluminum binary alloys and the earlier pioneering work of Galvele and his co-workers on conventional cast-aluminum binary alloys. D. Ternary and Commercial Alloys As discussed earlier, the corrosion behavior of Al–Li binary alloys in saline solutions is independent of alloy heat treatment other than when alloys are grossly overaged and δ,Al–Li, a highly anodic phase, forms at grain boundaries (149,150). Copper additions to Al–Li alloys significantly modifies this situation with thermal aging at temperatures around 170°C promoting copper-containing precipitation at subgrain boundaries, T1(Al2CuLi) followed by T2(Al6CuLi3) at high-angle grain boundaries when tempers approach peak strengths (168). The intergranular corrosion susceptibility for both Al–Li–Cu alloys and the commercial alloy 2090 have been explained by the selective dissolution of a copperdepleted zone (147), as previously proposed for Al–Cu alloys (137,142). The situation for Al–Li–Mg–Cu alloys (e.g., 8090) is slightly different, with the presence of magnesium leading to different precipitates being developed during aging. The influence of alloy temper on the mode of localized attack and pitting potentials for these alloys, including 2091 (169), however, closely resembles that Aluminum Alloys 231 Fig. 34 Effect of aging time at 170°C on the grain boundary and matrix pitting potentials for 2024 in a de-aerated 1M NaCl solution adjusted to 10. (Data from Ref. 143.) observed for the Al–Cu and the Al–Mg–Cu alloy systems, (compare Figs. 30, 34, and 35 for Al–3.3 Cu, 2024, and 8090, respectively). Similar trends have been reported for Al–Zn–Mg–Cu alloys by Maitra and English (170). They report that the temper dependence of the pitting potentials for the Al–Zn–Mg–Cu alloy 7075 in an aqueous saline solution is remarkably similar to that reported by Muller and Galvele (142) for Al–Cu alloys. Their anodic polarization studies on 7075 in a de-aerated 3.5% NaCl solution show the following: 1. A single pitting potential was detected for the solution heat-treated and rapidly quenched condition. 2. Two pitting potentials existed for the peak aged temper, the less noble one being associated with the local breakdown in the grain-boundary region and second corresponding to pitting of the matrix. 3. Sufficiently overaged alloy microstructures displayed a single pitting potential that corresponded to pitting of the copper-depleted solidsolution matrix. 232 Holroyd Fig. 35 Efect of aging time at 170°C on the grain boundary and matrix pitting potentials for 8090 in a de-aerated 3.5 (wt%) NaCl solution adjusted to 6.5. (Data from Ref. 148.) Although the microstructural details of the precipitates formed differ for the various age-hardening aluminum alloy systems, the influence of alloy temper on the modes of localized corrosion suffered are self-consistent for the Al–Cu(–Mg)-, Al–Zn–Mg–Cu-, and Al–Li-based systems (i.e., 2xxx, 7xxx, and 8xxx series alloys). VI. CONCLUDING REMARKS: ROLE OF SURFACE FILMS DURING THE CORROSION OF ALUMINUM ALLOYS Based on available published information, it is reasonable to conclude that the ‘‘oxide’’ layers existing on aluminum or its alloys will ‘‘chemically’’ react when exposured to aqueous environments and this will be a precursor to any subsequent corrosion process that may subsequently initiate. This hypothesis is compatible with previous suggestions for the initiation of localized corrosion because it is a ‘‘global’’ surface process occurring prior to any ‘‘localized’’ processes such as blister development (104–109, 152) and/ Aluminum Alloys 233 or the development of critical local solution chemistries at potential initiation sites for pitting (138,164,171), grain boundary (137,142), or crevice (172–175) corrosion. These latter processes are usually associated with some heterogeneous alloy microstructural feature at the metal–film interface alloy surface beneath the surface film or possibly associated with a flaw/defect in the oxide (81,82). The ‘‘chemical’’ reaction experienced by the ‘‘oxide’’ film is highly dependent on the environmental conditions to which it is exposed. The nature of the ‘‘oxide’’ and its thickness may influence these interactions with these effects usually being kinetic, although there are instances where the type of pre-existing ‘‘oxide’’ can determine whether corrosion initiates (11,12,15). When aluminum or aluminum alloys are exposed to aqueous environments (or water vapor), the ‘‘oxide’’ surface films are modified via hydration and/or ‘‘dissolution’’ processes. The extent and details of these processes are dictated by the ionic species present and the solution pH, temperature, pressure, and flow rate of the environment. For most instances when aluminum and its alloys are exposed to aqueous solutions with pH’s in the range 4–9, a porous hydroxide (often pseudo-boehmite and/or bayerite) forms by nucleation and growth from soluble aluminum species generated from the pre-existing ‘‘oxide’’ (29,39,73). The overall reaction is controlled by what happens to the dissolved aluminum species immediately after immersion. Typical situations include the following: 1. ‘‘Oxide’’ hydration with growth and densification of a hydroxide layer retarding reaction rates and eventually promoting self-stifling. A typical example of this is the filming of pure aluminum when exposed to distilled water, as described earlier in this chapter. 2. Corrosion inhibition by insoluble complexes forming on the oxide surface and preventing further ‘‘oxide’’ hydration [as suggested by Vermilyea and Vedder (39)]. A good example of this is the filming behavior reported for pure aluminum in dilute aqueous phosphate solutions (10⫺3M–10⫺2M), where the hydration of the oxide is thought to be completely suppressed with no evidence of oxide dissolution other than dissolved aluminum species interacting with phosphate species to form zero-charge complexes that are stable on the oxide and hinder the penetration of water molecules through the oxide (176). 3. ‘‘Oxide’’ hydration initiates as in situation 1 but restricted geometry (e.g., a tight crevice) promotes local concentration buildup of dissolved aluminum species causing local solution pH’s to become increasingly acidic. Depending on the anions present, the possibility that film ‘‘dissolution’’ and ‘‘modification’’ occurs in some circumstances can lead to the onset of localized corrosion. An excellent example of this is given in the detailed study of the crevice corrosion of pure aluminum 234 Holroyd in 1M NaCl conducted by Baumgartner and Kaecshe (175). These researchers showed the following: (a) Crevice corrosion can be triggered by unstable micropitting at potentials as low as around ⫺1.040 V (SCE), some 0.3 V below the pitting potential Ep and at least 0.1 V (SCE) below the repassivation potential ER for the 1M NaCl bulk environment. (b) Unstable micropitting occurred during the incubation period for crevice corrosion, which also occurs when aluminum surfaces in noncrevice situations are exposed to aluminum chloride or acidified NaCl solutions. (c) When a critical solution chemistry is established inside a crevice, the mode of attack switches to a more general type of attack, operating at a low current density of around 10 µA/cm2. Based on these findings, Baumgartner and Kaecshe (175) suggest that the buildup of a critical acidic electrolyte is a sufficient requirement for the initiation of crevice corrosion. Other researchers have disagreed, suggesting that the requirement is either a critical concentration of aluminum ions that is the chloride ion concentration dependent (174) or the presence of basic aluminum chloride (171). 4. Aggressive ionic species in the aqueous environment modify the ‘‘oxide’’ hydration process and gel-like anion-selective surface membranes form (90), allowing water entry and an inward migration of specific anions. In this scenario, dissolved aluminum species (mainly Al3⫹) and aggressive anions (e.g., Cl⫺) build up at the metal–film interface beneath the gel-like films and lead to the initiation of localized corrosion at active microstructural sites (Fig. 36). (It may be argued that situation 3 is a special case of situation 4) for which the local ‘‘bulk’’ environment inside a restricted geometry (e.g., a crevice or crack) changes and then process 4 occurs.) Recent time-lapse video studies (103) have shown that continuous gel-like corrosion films form on aluminum alloy surfaces exposed to saline environments. Gas bubble generation occurs at sites beneath these films and, in some cases, leads to local eruptions and the exposure of sites where localized corrosion has become established. Although it is clear that the mechanism of localized corrosion initiation needs further detailed study, it is reasonable to conclude that the initiation of localized corrosion for aluminum and its alloys in aqueous environments normally occurs by either situation 3 or 4, irrespective of whether it occurs under total immersion, thin-film, or wet–dry environmental conditions. Surface films may have several roles during the propagation phase of localized corrosion: Aluminum Alloys 235 Fig. 36 Schematic representation of a anion-selective gel-like surface film that has formed on an aluminum alloy surface and has facilitated local acidification and the possibility of localized corrosion initiation beneath the film (based on the idea presented by Sato (90)) (not to scale). 1. Preventing extensive lateral spread and the initiation of general corrosion. 2. Maintaining the occluded nature of the corrosion site by providing a total, partial, or intermittent seal-restricting migration, diffusion, or significant hydrodynamic flow into or out of the corrosion site. 3. A surface film coating the inside of a local corrosion site can effectively limit localized corrosion rates. The films fulfilling the above functions will tend to be different. An example of this is shown schematically in Fig. 37, in which the surface films formed on the surfaces of an aluminum alloy suffering localized corrosion in an aqueous saline environments are envisaged to change as a function of the depth of a localized corrosion site. Some experimental justification for this proposal exists (177, 178): 236 Holroyd Fig. 37 Typical surface films and the solution is pH’s developed as a function of depth within restricted geometries for aluminum and its alloys exposed to saline solutions. 1. Le and Foley (177) measured the solution pH within stress corrosion cracks for 7075-T651 in a saline solution and reported that the solution’s pH’s decreased as a function of crack length from 4.6 to 5.0 in the precracked region, through a region with a pH of 4.2–4.6 with evidence of corrosion product to the advancing crack-tip region with a pH of 3.0–3.2. 2. Holroyd et al. (178) reported that the solution’s pH’s within artificial crevices for aluminum alloys exposed to saline solutions vary with both time and crevice depth (Fig. 37) and that the measured pH’s are consistent those known for the aluminum oxychloride products (Table 4). Table 4 Saturation Concentrations and Solution pH’s (at 90% Saturation) for Aqueous Solution of Aluminium Oxychloride Salts (179) and Aluminum Chloride (180) Aluminum oxychloride salt Al(OH)Cl2 Al(OH)2Cl Al2(OH)5Cl AlCl3 Saturation concentration (M) Solution pH (90% sat.) 3.02 3.12 2.34 3.11 2.85 2.94 3.67 0.15 Aluminum Alloys 237 Further support for these ideas is provided by Wong and Alkire (179), who have characterized solution chemistries developing within naturally occurring 100µm-deep pits in 99.999% purity aluminum in 1M NaCl. Their results indicate the following: 1. The solution’s pH’s in pits are between 3 and 4 and the acidity can be explained by the hydrolysis of Al3⫹ ions, as has been suggested by others for solutions within pits (181), crevices (87,182,183), and cracks (87,183–188). 2. Eighty-four percent of the measured dissolve aluminum in the pits is present as monomeric species and the remainder is present as polymeric species with no evidence of solid gels. 3. The hydrolyzed aluminum species react within the pit environment to form basic salts that have nuclear magnetic resonance spectra very similar to those given by synthetic solutions of Al(OH)Cl2 or Al(OH)2Cl. 4. Current densities from single pits decreased roughly with the square root of time and are consistent with the corrosion rate being controlled by the rate of dissolution of an aluminum oxychloride salt film formed by the precipitation of dissolution products. Several authors (181,185,186,189,190) in the past have suggested that the solution’s pH’s developed within restricted geometries (e.g., pits, crevices, and cracks) for aluminum alloys exposed to aqueous saline environments are controlled by acid hydrolysis involving Al(OH)3. For instance, Pryor (181) suggested that in practical situations, the solution’s pH in aluminum pits would not fall below 3.5–4.0 and he calculated a minimum possible pH of 2.8 using thermodynamic data for the chemical equilibrium process: Al3⫹ ⫹ 3OH⫺ ↔ Al(OH)3 (2) A major flaw in the above argument is that the species Al(OH)3 is unstable in solutions with pH’s below 5 (87,177,191,192) and, therefore, it is unreasonable to assume that pH’s are being controlled by Eq. (2). Recent theoretical work by Galvele (192) indicates that no solid-reaction products are expected to form in aluminum crevice solutions with pH’s of 3 and below, irrespective of bulk solution pH’s, and only 10% of the products would be solid in a crevice solution with a pH of 4. Several researchers have suggested that the solution’s pH’s within restricted geometries are well explained by the hydrolysis of Al3⫹ ions to form AlOH2⫹ and H⫹ ions via the equilibrium process (62,87,177,178): Al3⫹ ⫹ H2O ↔ AlOH2⫹ ⫹ H⫹ (3) The more recent work of Wong and Alkire (179) is consistent with Eq. (3) but indicates the difficulty in differentiating between Al(OH)2Cl and Al(OH)Cl2 and 238 Holroyd determining which is dominant within restricted geometries. One possibility is that neither is dominant and their balance within an occluded cell reflects the local solution’s pH. An example of this is shown schematically in Fig. 37 and a qualitative description of the chemical equilibrium processes controlling the local solution and film-formation processes within a restricted geometry for an aluminum alloy exposed to an aqueous solution is given by Al3⫹ ⫹ H2O ↔ AlOH2⫹ ⫹ H⫹ ↓ ↓ ⫹Cl⫺ ⫹Cl⫺ ↓ ↓ AlCl2⫹ Al(OH)Cl⫹ Al(OH)2Cl2⫺ In view of the experimental evidence now available, the previous suggestions that the presence of AlCl3 is needed to activate passive aluminum (171) or saturated AlCl3 must form at the bottom of pits for them to remain active (126,171,193) are clearly inappropriate. Perhaps the most direct evidence for this is the excellent agreement between the measured solution chemistries within restricted geometries [e.g., dissolved aluminum concentrations, the solution’s pH’s and the NMR (nuclear magnetic resonance spectroscopy) data] and those measured for species other than AlCl3. If it is accepted that saturated AlCl3 does not form at the base of restricted geometries, we need to modify the data we input into the Galvele pit model earlier in this chapter [see Eq. (1) and Table 3] so that the pitting potential, E *p , reflects the actual environmental conditions developed within a pitlike geometry rather than that developed when an aluminum alloy is exposed to a saturated aqueous solution of AlCl3 for 2 h (140). Typical published data for electrochemical and environmental conditions developed within aluminum crevices and cracks are given in Table 5. If we assume that the potentials developing at the base of active crevices or propagating stress corrosion cracks are realistic values for E p* and substitute into Eq. (1), Ep ⫽ E p* ⫹ η ⫹ ϕ ⫹ Einh (1) along with published pitting potentials and taking Einh ⫽ 0, we can calculate values for η ⫹ ϕ where η is the overpotential needed to draw a net anodic current through a restricted geometry and ϕ is the potential gradient within the restricted geometry. Experimental work is needed to validate the (η ⫹ ϕ)calc values given Table 6. Vermilyea and Vedder (39) studied the filming behavior of pure aluminum Aluminum Alloys Table 5 Environmental Conditions Developed Within Restricted Geometries for Various Aluminum Alloys Restricted geometry Alloy Solution pH Internal potential External potential Al3⫹ conc. (M) Cl⫺ conc. (M) Bulk NaCl conc. (M) Ref. Crevices 1 mm diameter 10 mm deep 80 mm deep, 90 µm wide 15 mm diameter ⫻ 30 mm Cracks 6 mm long 15 mm long 20 mm long Pit 100 µm deep 1199 2017 3003 6063 7475-T6 7475-T6 LC4-T6 7075-T6 4 4 3.5 3.5 4 7475-T6 LC4-T6 3.4 3–3.2 2.7 3.2 S.P. Al 3–4 ⫺0.830 ⫺0.825 ⫺0.844 ⫺0.916 ⫺0.940 ⫺0.870 ⫺0.920 ⫺0.925 ⫺0.780 ⫺0.773 ⫺0.890 ⫺0.780 ⫺0.790 ⫺0.830 ⫺0.850 ⫺0.832 ⫺0.830 0.003 0.25 0.025–0.37 0.1 0.75 2–3 1.82 1.0 0.5 0.5 0.5 0.6 0.51 0.6 1.0 1.0 0.51 0.6 187 182 182 182 194 87 184 187 177 87 184 179 239 240 Holroyd Table 6 Calculated Values for (η ⫹ ϕ) from Eq. (1) Using Published Data Alloy Pure Al (in 1M NaCl) 3003 (in 0.5M NaCl) 7475-T6 (in 0.6M NaCl) Ep (Ref.) E*p (Ref.) (η ⫹ ϕ)calc. ⫺0.740 (222) ⫺0.830 (187) crevice ⫺0.844 (182) crevice ⫺0.870 (194) crevice 0.090 ⫺0.713 (222) ⫺0.760 (170) 0.131 0.110 in boiling water containing 10⫺3M of various anions. Takahashi et al. (176) have classified these results into four groups as a function of their ability to inhibit hydration: 1. No inhibition: borate, MnO4⫺, ClO4⫺, NO3⫺, NO2⫺, Cl⫺, and CO32⫺ 2. Moderate inhibition: IO3⫺, SO42⫺, SeO42, GeO42⫺, CrO42⫺, citrate, oxalate, and MNO42⫺ 3. Strong inhibition: Teo42⫺, SiO42⫺, WO42⫺, AsO42⫺, and IO42⫺ 4. Extremely strong inhibition: PO43⫺ Generalization of these findings needs care because the behavior of some anions are sensitive to their concentration and/or pH and temperature. Phosphate ions is a good example. At low concentrations (10⫺3M–10⫺2M), oxide hydration is completely suppressed (as discussed earlier in this chapter); at intermediate concentrations (10⫺1M), oxide films initially suffer both dissolution and hydration with a ‘‘dissolution–deposition’’ process generating a surface film; at higher phosphate concentrations, the dissolution of the oxide is accelerated by the formation of anion complexes that themselves hydrolyze to form Al-hydroxo-phosphate complexes that then undergo polymerization to deposit ‘‘basic phosphate’’ or ‘‘hydrated oxides containing phosphate’’ layers (176). The presence of several anions in an aqueous solution may promote unexpected results with regard to the corrosion performance of aluminum alloys in aqueous environments. Three examples follow. Becerra and Darby (195) found that the corrosion rates for AA1100-H14, AA5052-H32, and AA6063-T4 sheet materials all increased with additions of sodium bicarbonate (0–225 ppm) or cupric sulfate (0–2 ppm) added as single additions to a 3.5 wt% NaCl solution. The increased corrosion rates were greater than simply additive when the additions were added simultaneously. The synergistic effect was attributed to copper plating out on the aluminum alloy surface and acting as a cathode [as suggested by Davies (196) and Fraker and Ruff (197)] and the bicarbonate’s buffering capability limiting the local solution’s pH Aluminum Alloys 241 changes at the cathodes and reducing the tendency for local protective film formation. Overaged Al–Zn–Mg–Cu alloys usually provide a good resistance to intergranular corrosion in aqueous saline solutions (170). Maitra and English (198) have shown that small additions of nitrate and sulfate ions (particularly nitrate ions) can cause 7075-T7351 to suffer intergranular corrosion during laboratory testing. They also found that the relative susceptibilities promoted by the different solution chemistries were reflected by the difference between the grain-boundary and matrix pitting potentials. These results suggested the possibility that overaged high-strength Al–Zn–Mg–Cu alloys, although being highly resistant in a saline environments containing just chloride ions, may be prone to either intergranular corrosion or stress-corrosion cracking (SCC) in certain industrial environments (198). To date, there is no evidence of service performance issues with localized corrosion occurring in overaged high-strength Al–Zn–Mg–Cu alloys. However, it is interesting to note that the stress-corrosion performance of 7075-T7551 and 7075-T7351 during a extensive ongoing test program on various aluminum alloys in rural, industrial, and seacoast environments (199) has identified that these alloy tempers are significantly more susceptible to SCC in an industrial environment. Yet another example of synergetic effects between anions influencing the localized corrosion behavior of aluminum alloys is that reported for the Al–Li– Mg–Cu alloy (8090) and the Al–Li–Cu alloy (2091) where small additions of sodium sulfate to aqueous sodium chloride solutions have had marked effects on the alloy’s SCC susceptibility (200,201). For 8090-T651, Craig et al. (200) demonstrated that whereas 8090-T6 was resistant to SCC when under totalimmersion conditions in an aqueous 0.6M sodium chloride solution, it became susceptible with small additions of sulfate ions but highly resistant with higher levels of sulfate additions (Fig. 38). Similar findings have been reported by Marsac et al. (201) for 2091, albeit this alloy simultaneously suffers a susceptibility to intergranular corrosion that seems to be independent of the solution’s sulfate concentration. A likely explanation of the sulfate effect at low concentrations is that the surface filming process is modified by the introduction of a thermodynamically more attractive reaction path allowing the formation of aluminum oxysulfate species which will form in competition with the aluminum oxychloride species, as has been described by Foley and Nguyan (202). For the high concentrations of sulfate ions, it could be that basic aluminum sulfate films form. Based on their findings, Craig et al. (200) predicted that 8090-T651 would suffer SCC when under total-immersion conditions in artificial or natural seawater. This was observed, and the susceptibility being lower than predicted on the basis of the solution’s sulfate-to-chloride ratio (Fig. 38) was attributed to filming effects promoted by the additional soluble additives in seawater (e.g., magnesium salts). 242 Holroyd Fig. 38 Time to failure for peak aged 8090 stress-corrosion samples tested under totalimmersion conditions in a range of aqueous environments containing various concentrations of sodium sulfate and sodium chloride at 20°C. (Data from Ref. 200.) A. Hydrostatic Pressure Effects To the author’s knowledge, the available published literature on the effects of hydrostatic pressure on the films formed during the localized corrosion of aluminum alloys is limited. Although most laboratory (203–205) and field testing (206–208) studies indicate that the pitting susceptibilities of aluminum and its alloys exposed to seawater generally increase with increasing hydrostatic pressure, data have been published showing that corrosion rates may be insensitive or actually decrease with increasing pressure (209,210). In the late 1960s, the increased pitting susceptibilities were attributed to hydrostatic pressure increasing chloride ion activity (211) and its penetration into passive layers. It is now realized that the surface films are important and, in some instances, can control corrosion performance (205,210). A good example of the importance of surface films is provided by Beccaria and co-workers (205,210), results showing that increasing hydrostatic pressure increased the corrosion rate for pure aluminum exposed to 10⫺2 M Na2SO4 (205) but decreased corrosion rates for 6061-T6 exposed to seawater (210). Aluminum Alloys 243 In the case of the pure aluminum exposed to aqueous Na2SO4 solutions, Beccaria and Poggi (205) found that sulfate ion additions in the range 10⫺2M– 10⫺1M reduced corrosion rates for all the hydrostatic pressures tested in the range of 1–300 atm. Corrosion rates in solutions containing 10⫺2M sulfate increased with increasing hydrostatic pressure, whereas corrosion rates in solutions containing 10⫺1M sulfate were insensitive to hydrostatic pressure. The observed behavior at the lower sulfate concentration was deemed independent of sulfate ions and associated with hydrostatic pressure promoting the incorporation of aluminum hydroxide into surface films that would be more soluble at higher pressures and therefore less protective. For the higher sulfate concentration, the authors suggested that aluminum sulfates formed in competition with the hydroxides and the net effect was that corrosion rates were insensitive to increasing pressure. Experimental results reported by Beccaria et al. (210) showing that the corrosion rates for AA6061-T6 exposed to seawater decreased with increasing hydrostatic pressure would seem to be contrary to their previous work and that reported by others. Their explanation, based on surface film analysis using Xray photospectroscopy and infrared spectroscopy, was that hydrostatic pressure promotes a more protective and less hydrated surface film containing magnesium oxide and pseudo-boehmite, as opposed to a less protective hydrated bayerite [Al(OH)3] film that forms at atmospheric pressure. These findings are extremely interesting and deserve further evaluation because they suggest that hydrostatic pressure can induce significant changes to the surface films forming on aluminum alloys in aqueous solutions. B. Influence of Flowing Conditions Erosion corrosion of aluminum alloys in aqueous environments due to flowing conditions usually become evident when flow rates exceed 3 m/s and the loss of metal increases sharply with increasing flow rates in excess of about 9 m/s (212,213). Localized corrosion is sensitive to solution flow rates below 3 m/s, with evidence of low flow rates up to around 1.5 m/s being beneficial (214,215) or at least not detrimental, and higher flow rates up to 3 m/s having minor or a gradually increasing effect on corrosion rates (212,213). These effects may differ depending on whether or not localized corrosion has initiated and is well established. In the case of corrosion initiation, low flow rates may lead to surface films that prevent or delay initiation. For localized corrosion that is well established and in a propagation mode, the influence of external solution flow is dependent on whether solution chemistries are influenced within the local corrosion sites. As with the hydrostatic pressure effects, the available literature on the effect of flowing aqueous conditions on the localized corrosion of aluminum and its alloys is limited (215,216). In theory, mathematical modeling is feasible for the 244 Holroyd flowing conditions, despite the combination of localized corrosion, general corrosion, and bulk environmental flow, creating a complex set of differential equations that are difficult to solve (217). Most published studies on the effect of flowing conditions on the corrosion of aluminum alloys have involved marine environments (212–216,218–220) and are associated with applications such as marine vessels, structural and nonstructural uses on offshore oil platforms, ocean thermal energy conversion (OTEC) heat exchangers, and desalination plants. A notable exception being Witt’s (221) studies on the effects of flowing conditions have on aluminum alloy (AA6063 type) domestic radiator systems. Godard and Booth (218) reported that the pitting corrosion rates for seawater flowing through 1S-H18 aluminum (0.02 Cu, 0.21 Fe, 0.11 Si, 0.016 Ti) pipes initially decreased with increasing flow rates and then increased with further increases in flow rate. Unfortunately, only relative flow rates are available. Beneficial effects of low flow rates on localized corrosion rates have also been reported for various alloys used in OTEC heat exchangers (214,215) and in domestic hotwater radiator systems (221). Larson-Basse and co-workers (214–216,219) have evaluated the corrosion performance of a range of aluminum alloys [AA3003 and 5052 and Alcald (7072) 3003 and 3004] in flowing seawater in OTEC heat exchangers using flow rates in the range 1.4–2.4 m/s. The observed localized corrosion behavior in the tropical surface water differs significantly from that promoted in the cold, deep ocean water (214,216). In the warm surface water, a surface film forms after an initial short period (5–10 days) of relatively rapid corrosion which is comprised of both scale minerals precipitated from the seawater and aluminum corrosion products that are alloy dependent (214). This film is protective the corrosion rates fall to around 3 µm/year for all the alloys tested. Unfortunately, the film thickness increases in a parabolic manner and is sufficiently detrimental to the systems heattransfer characteristics that it has to be regularly removed. This introduces challenging problems with both material selection and choice of method to remove scales (216). Exposure of aluminum alloys to the cold, deep ocean water under flowing conditions as used in OTEC heat exchangers introduces little or no tendency for biofouling; however, the alloys usually suffer localized pitting, with the incubation times being sensitive to small variations in solution chemistry (214) and pitting tendencies increasing significantly with decreasing flow rates (214). Witt (221) has reported that residual water left in a hot-water radiator system for a few hours after a successfully operating system has been emptied can locally disrupt the protective film that had previously provided protection, and intergranular corrosion can result during the subsequent operation of the radiator system. Once again, this provides clear evidence that the corrosion performance of an aluminum alloy in a given aqueous environment is not always predictable Aluminum Alloys 245 and can be dictated by the surface film already present or one forming upon immersion in an aqueous environment. 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CA Witt, Aluminium 56:398, 1980. 8 Magnesium Alloys Mike J. Danielson Pacific Northwest National Laboratory, Richland, Washington I. INTRODUCTION Aluminum and zinc are the two major alloying elements added to magnesium to increase its strength, ductility, and resistance to general corrosion attack. Magnesium alloys are of engineering interest because of their low density and high strength-to-weight ratio. Magnesium alloys are about two-thirds the density of aluminum alloys and about 25% lighter than aluminum alloys with similar stiffness. Other advantages are good machinability and castability into complex configurations. Magnesium alloys are used in applications where strength and lightness are at a premium, such as aircraft, space applications, automotive applications, and certain consumer uses (i.e., luggage, computer frames, audio equipment, lawnmowers). In 1990, most alloy use was in die castings involving one composition, AZ91D (1). Under dry or low-humidity applications, these alloys have excellent materials properties, but when immersed or at high relative humidity, they can suffer severe localized corrosion attack, stress-corrosion cracking, and corrosion fatigue. There have been a number of reviews (2–6) on the materials performance of magnesium alloys. This review will focus on the effect of the environment on the materials performance, including general and localized corrosion attack, stress-corrosion cracking, and corrosion fatigue under a variety of relevant engineering environments. This review will (a) summarize the early work and (b) cover the more recent work since 1993, the most recent review. The intent will be (a) to describe the underlaying principles that control the behavior of these alloys, (b) describe the failure modes and their character, (c) describe methods 253 254 Danielson that mitigate the failure processes, (d) direct the engineer in searching and evaluating the literature, and (e) assist the engineer in carrying out new testing in relevant environments. The literature is quite complete on the subjects of generalized and localized attack but very sparse on stress-corrosion cracking and corrosion fatigue. On these latter two topics, benchmark papers are pointed out that will be used to reveal the general character of magnesium alloys. After reading this chapter, the engineer will have a good grounding in using these alloys in practical applications and avoiding corrosion-related failures. II. NOMENCLATURE Successful use of these alloys requires an understanding of how the alloys are identified. The American Society for Testing and Materials (ASTM) created a standard system for naming magnesium alloys and tempers which has been in use since 1952 (7). This system is clearly described by Hough et al. (8), and an abbreviated version of their original table is shown in Table 1. The system is most easily learned by following an example while referring to Table 1, using AZ91E-T6 as the magnesium alloy. The letters AZ refer to the two alloying elments in highest concentration, aluminum and zinc (highest concentration shown first). The next two numbers, 91, refer to the approximate amount of each alloying element, namely 9 wt% Al and 1 wt% Zn, rounded up to the largest whole number. The third part of the description, E, is an additional designator of the alloy and may have importance in specifying additional compositional details. The fourth part, T6, refers to the temper, which in this case is a solution anneal followed by some artificial aging. Table 2 shows the composition and yield strength of some commonly used cast magnesium alloys. III. CORROSION BEHAVIOR A. Basic Metallurgy of Magnesium Alloys Affecting the Materials Performance Aluminum and zinc are the two elements most often used to alloy with magnesium to improve its strength and corrosion properties, and their phase diagrams are shown in Figs. 1 and 2, respectively (9). Generally, the aluminum content does not exceed about 10 wt% or the zinc content about 6 wt% in commercial alloys. This is because the solubility of aluminum in magnesium is 12.8 wt% at 437°C, and exceeding this concentration will result in the formation of the βphase, Mg17Al12. However, the solid solubility of Al significantly decreases with temperature (1.5 wt% at ambient), and some Mg17Al12 will usually be present depending on the quenching rate and tempering. The solubility of zinc in magne- Magnesium Alloys 255 Table 1 ASTM Identification System for Magnesium Alloys and Temper First part Indicates the two principal alloying elements using two coded letters: A ⫽ B ⫽ C ⫽ D ⫽ E ⫽ F ⫽ G ⫽ H ⫽ K ⫽ L ⫽ M⫽ N ⫽ P ⫽ Q ⫽ R ⫽ S ⫽ T ⫽ W⫽ Y ⫽ Z ⫽ aluminum bismuth copper cadmium rare earth iron magnesium thorium zirconium lithium manganese nickel lead silver chromium silicon tin yttrium antimony zinc Second part Third part Two numbers indicating the weight percent (wt%) the amounts of the two principal elements A letter designation that distinguishes between two alloys with the same composition of the two principal elements Fourth part Temper designation: F ⫽ as fabricated O ⫽ annealed H10, H11 ⫽ slightly strain hardened H23, H24, H26 ⫽ strain hardened and partially annealed T4 ⫽ solution heat treated T5 ⫽ artificially aged, only T6 ⫽ solution heat treated and artificially aged T8 ⫽ solution heat treated, cold worked, and artificially aged Source: Refs. 7 and 8. sium is 6.2 wt% at 340°C, and exceeding this concentration will result in the formation of a variety of complex intermetallics, such as MgZn2. Because the solubility of zinc also decreases with temperature (⬃1.7 wt% at 150°C), zincrich intermetallics may also be present. Mn and, more recently, Zr are added to the alloy to act as grain refiners to improve the strength. Due to their fairly low solubility, they will also be present as intermetallics in the matrix. Fe, Ni, and Cu are very insoluble impurities, which above a certain threshold concentration are well known to have extremely detrimental effects on the localized corrosion behavior. They are present as intermetallics which act as cathodic depolarizers in the creation of a localized microgalvanic cell. Magnesium alloys make use of a number of other minor metal additions of low solubility such as the rare earths, 256 Table 2 Nominal Compositions and Yield Strengths of Selected Magnesium Cast Alloys Alloy Mn Zn 10.1 6.0 7.6 8.7 8.7 9.0 0.1 0.15 0.13 0.13 0.13 0.10 3.0 0.7 0.7 0.7 2.0 2.7 3.3 3.3 1.0 0.25–0.75 6.0 4.2 5.8 5.7 4.6 6.0 Other Th 0.7 0.6 0.7 0.7 0.7 0.7 0.7 0.7 0.7 2.7 0.7 0.7 Zr, Zr, Zr Zr Zr Zr, Zr, Zr, Zr, Cu Zr, Zr, 1.8 0.7 Zr 0.7 Zr Note: In most cases, Cu (⬍0.08%), Ni (⬍0.01%), and Fe (⬍0.005%) are kept at very low values. Source: Refs. 7 and 8. 1.5 Ag, 2.1 Di 3.3 RE 2.5 2.5 4.0 5.2 Ag, 2.1 Di Ag, 1.0 Di Y, 3.4 RE Y, 3.00 RE 1.2 RE 2.6 RE Yield strength (MPa) 150 130 83 145 145 150 195 110 105 90 55 195 205 165 172 125 140 190 170 165 195 Danielson AM100A-T61 AZ63A-T6 AZ81A-T4 AZ91C AZ91E-T6 AZ92A-T6 EQ21A-T6 EZ33A-T5 HK31A-T6 HZ32A-T5 K1A QE22A-T6 QH21A-T6 WE43A-T6 WE54A-T6 ZC63A-T6 ZE41A-T5 ZE63A-T6 ZH62A-T5 ZK51A-T5 ZK61A-T6 Al Magnesium Alloys 257 Fig. 1 Phase diagram of the Al–Mg system. which improve certain mechanical properties but also may form precipitates (contribute to the strengthening by precipitation hardening) within the matrix. An article by Beldjoudi et al. (10) examined the microstructure of Mg–9 Al, AZ91, and Mg–3 Al alloys. In the T4 (solution annealed) condition, the alloys were all single phase (dendritic substructure) with the Al uniformly distributed in the Mg matrix, there being no evidence of segregation. Precipitates of the type AlMnFe (globular) and Mg2Si (polygonal) were distributed uniformly throughout the matrix. Aging the Mg–9 Al alloy to the T6 condition resulted in the decrease of solid-solution Al from 9% to 3% and the formation of a lamellar β-Mg17Al12 at the grain boundaries. The T6 condition did not affect the size and distribution of the AlMnFe-containing precipitates, but the Mg2Si were now present primarily along grain boundaries. Zn did not affect the microstructure. The observations of Lunder (11,12) on AZ91 in the F, T4, and T6 tempers reinforce those of Beldjoudi, with the additional observation that the Zn is always in solid solution. The as-cast condition (F) contained a great deal of β-Mg17Al12 at the grain boundaries as well as being distributed throughout the matrix, and the solid-solution Al was distributed nonuniformly. The intermetallics, AlMnFe and Mg2Si, were present throughout the matrix and did not appear to be altered by heat treatment. 258 Fig. 2 Danielson Phase diagram of the Mg–Zn system. In summary, the presence of a variety of intermetallics in the matrix, the grain boundaries, and at other inhomogeneities has a profound affect on the corrosion performance of this entire class of alloys. Heat treatment can change the distribution and microchemistry of the intermetallics and it should be expected that the corrosion degradation processes will be a strongly dependent on the heat treatment. The detailed interaction between the intermetallics and the degradation mechanisms has not been adequately studied to date. B. Basic Electrochemistry of Magnesium Alloys Affecting the Materials Performance Magnesium is one of the most anodic (reactive) metals in aqueous solution, being next to the very reactive metals Li and Na in the electromotive series. A Pourbaix diagram (13) is shown in Fig. 3 that demonstrates that the electrochemical potential for the reaction Mg 2⫹ ⫹ 2e⫺ → Mg 0 Magnesium Alloys 259 Fig. 3 Pourbaix diagram for the Mg–H 2 O system. is more than 2 V below the hydrogen evolution ‘‘a’’ line, indicating that there is a large thermodynamic driving force for the spontaneous reaction (corrosion) of the Mg with water with the simultaneous evolution of hydrogen gas. Pourbaix diagrams only show the thermodynamic propensity for reaction; the actual kinetic rate must be determined by a direct physical measurement. Figure 4 shows the measured open-circuit corrosion potential of various Al–magnesium alloys in a common testing environment, 3.5 wt% NaCl. The significant feature of this plot is that the potential (written as a reduction potential) becomes more negative as the magnesium content increases, indicating that the driving force for the corrosion reaction increases with Mg content. Because many of the alloying agents 260 Danielson Fig. 4 Effect of the magnesium composition in Al–Mg alloys on the open-circuit corrosion potential exposed to 3.5% NaCl. have limited solubility in the magnesium matrix, intermetallics and impurities will be commonly present in the matrix as precipitates. For example, β-Mg17 Al 12 (56 wt% Mg) will be present near grain boundaries or other precipitates, and because its corrosion potential is less negative than the surrounding Mg matrix (see Fig. 4), it will be cathodically protected by the adjacent matrix that is richer in magnesium. One of the unusual features of magnesium alloys is that these precipitates are cathodic relative to the matrix; consequently, they act as sites for hydrogen evolution. This results in the creation of microgalvanic cells that are in intimate contact with each other. It appears that the corrosion rate of magnesium alloys is controlled by the rate of the hydrogen evolution reaction, and often these intermetallic cathodic sites have greater hydrogen evolution kinetic rates than the magnesium matrix itself. This explanation was used early in the development of Mg alloys to explain why the Fe, Cu, and Ni intermetallics, which support unusually large hydrogen evolution rates, are worse than any others for driving the galvanic process (14). Recently, Lunder et al. (15) synthesized a number of the normally present intermetallics [e.g., Al 6 Mn, Al 6 Mn(Fe), Al 3 Fe, β-Mg 17Al 12, Mg 2Si, and Al 4MM, where MM ⫽ Misch metal] and (a) measured their opencircuit corrosion potential in pH 10.5, 5% NaCl solution saturated with Mg(OH) 2 and (b) determined the hydrogen evolution rate when polarized to ⫺1.6 V (saturated calomel reference electrode scale), the open-circuit potential of magnesium. The hydrogen evolution rate was greatest for the Fe-containing intermetallics and least for Al 6Mn and Mg 2 Si. The β-Mg17Al12 gave intermediate values between these two extremes. Clearly, the composition of the intermetallics can have a very important role in controlling the localized corrosion performance, and the alloy purity is the starting point for controlling the intermetallic composition and improving the corrosion behavior of Mg alloys. Because the precipitates are less Magnesium Alloys 261 reactive than the matrix, an unusual type of localized attack occurs: the precipitates are protected and attack starts adjacent to them in the matrix. Careful microstructural evaluation is needed to reveal this type of morphology. Magnesium corrodes in neutral and basic water to form a magnesium hydroxide film that can be very protective under certain specific conditions: Mg 0 ⫹ 2H 2O → Mg(OH) 2 ⫹ H 2 (g) The pH and certain anions, such as chloride, can damage the protective film, leading to large increases in the corrosion rate. The composition of the film is a function of the alloy composition and the impurities in the environment (16,17). In summary, magnesium alloys are vulnerable to corrosive attack because of the extreme reactivity of the magnesium within the alloy and the intimate mixture of intermetallics that can act as localized cathodes, creating microgalvanic cells. The composition of the intermetallics can strongly affect the corrosion performance of the magnesium alloy. C. Localized and General Corrosive Attack (Aqueous Solutions) Magnesium alloys are prone to attack in aqueous solutions, particularly saltwater. The presence of certain metallic impurity phases that makes the alloys vulnerable has been recognized for over 50 years. Hannawalt et al. (14) carried out a pioneering study in 1942 that defined the ‘‘tolerance limits’’ for a number of the common impurities. Their effects on the corrosion rate in saltwater are shown in Fig. 5. Basically, Na, Si, Pb, Sn, Mn, and Al have negligible effects up to 5 wt% and Fe, Ni, Co, and Cu have extremely detrimental effects at very low concentrations. These tolerance limits are approximately 0.017, 0.1, and 0.004 wt% for Fe, Cu, and Ni, respectively, and these elements must be maintained at very low concentrations in order for magnesium alloys to have adequate corrosion resistance. Tolerance limits (3) are shown for some modern cast alloys in Table 3. Manganese is commonly added to commercial magnesium alloys at the nominal 0.2% level to act as a grain refiner and to increase the tolerance for Fe. Recent work indicates that the corrosion resistance is better correlated to the Fe/Mn ratio than to the Fe content (18). Manganese is believed to act by (a) forming AlMnFe intermetallic particles that precipitate out and fall to the bottom of the crucible, thereby reducing the iron content, and (b) the fact that the AlMnFe intermetallics remaining within the matrix are less efficient as cathodes as compared to metallic Fe, reducing the driving force for the galvanic cell. Table 4 (19) is included in order to benchmark the corrosion rate of a common magnesium alloy, AZ91, as a function of purity (Fe content) and heat treatment under ASTM B-117 salt spray conditions. Clearly, the corrosion rate of the commercial purity alloy (C suffix) is unacceptable in saltwater conditions in any heat-treatment condition, 262 Fig. 5 Danielson Corrosion rate of magnesium alloys as a function of various constituents. being nominally 15 mm/year. The high-purity alloy (E suffix) has a dramatically lower corrosion rate, and this reinforces the principle that improved corrosion resistance must come from improved purity, there having never been found any alloying elements that greatly inhibit the corrosion of magnesium. The results in Table 4 clearly show that heat treatment affects the localized corrosion behavior of the alloys. For instance, aging (T6 condition) results in the precipitation of Table 3 Tolerance Limits for Cast Alloys (ppm) Element Fe (with 0.2% Mn) Ni Cu AZ91 AM60 AS41 AE42 64 50 400 42 30 100 20 40 200 40 200 1000 Magnesium Alloys Table 4 Alloy AZ91C AZ91C AZ91E AZ91E AZ91E AZ91E Corrosion Behavior of AZ91 in Saltwater Grain size (µm) 187 66 146 78 160 73 Temper and corrosion rate (mm/year) Mn% Ratio Fe/Mn F T4 T6 T5 0.18 0.16 0.23 0.26 0.33 0.35 0.087 0.099 0.008 0.008 0.004 0.004 18 17 0.64 2.2 0.35 0.72 15 18 4 1.7 3 0.82 15 15 0.15 0.12 0.22 0.1 — — 0.12 0.12 0.12 0.1 263 264 Danielson the β-Mg 17Al 12 phase at the grain boundaries which are cathodic to the matrix, and yet this heat treatment results in the most beneficial condition for decreasing localized attack. The solution anneal case (T4) shows an anomolously high corrosion rate relative to the as-cast (F) and aged conditions (T6) (12,19) that is not satisfactorily explained (the solution annealed alloy should be more homogeneous that the cast or aged material). Beldjoudi et al. (10) reproduced these effects of heat treatment and clearly showed by electrochemical measurements that βMg 17Al 12 is less reactive than the matrix. Lunder et al. (11) examined the attack morphology on AZ91 in saltwater and found the pitting attack initiated adjacent to the intermetallic particles consisting of type AlMnFe, β-Mg 17Al 12 , or Mg 2Si. Attack followed the dendritic arms in the as-cast material, whereas filiform corrosion was observed for the T4 condition. Other alloying agents (Zr, Y, and rare earths Nd, Gd, Tb, Dy) are used to increase the strength of magnesium alloys, and they appear to have a small effect on the corrosion resistance as reported by King et al. (20) and Kamado et al. (21) in salt spray tests. King et al. (20) report corrosion rates for WE54 and WE45 of about 0.7 mg/cm 2 /day (1.4 mm/year), which are still substantial. Lunder et al. (12) determined that rare earths and silicon additions have a beneficial effect on the corrosion resistance, whereas zinc has a negligible effect. Magnesium alloys are attacked by acidic solutions but are quite resistant to alkaline environments and many hydrocarbon environments. Two exceptions to the acidic generalization are HF and chromic acid solutions, in which the alloy can be quite resistant (2,3). The alloys are very resistant to deionized water but will pit in potable water, but the threshold halide concentration for pitting is not defined in the literature. Except for fluoride, any of the halides, sulfates, and nitrates are damaging. Solutions containing dissolved Cu, Ni, or Fe ions are especially damaging because the ions will be reduced onto the alloy surface and act analogously to what happens when they are present within the alloy as impurities. Froats et al. (22) have a useful chart of organic and inorganic environments with the corrosion response of the magnesium alloys. Many of the results are prefaced with the excellent advice that the material should also be tested in the actual engineering environment to ensure a satisfactory result. Hallopeau et al. (23) examined the behavior of certain traditional inhibitor anions (SiO 32⫺ , PO 43⫺ , CrO 42⫺ , MoO 42⫺ ) on a Mg–9 Al alloy in a 0.5M sodium sulfate solution. The corrosion behavior was determined by using potentiodynamic and polarization resistance methods, which only evaluate the short-term performance. As expected, the Mg was selectively dissolved from the alloy, leading to a surface film enriched in Al. Some inhibition was observed in alkaline solutions, with SiO 32⫺ and CrO 42⫺ having an inhibiting effect in neutral solutions. No effect was observed for MoO 42⫺ and PO 43⫺. The effect of increasing the temperature is to increase the severity of attack. As an example, the corrosion rate, which is negligible in deionized water at ambi- Magnesium Alloys 265 ent temperature, increases to about 0.5 mm/year for the AZ alloys at 100°C (16). Clearly, magnesium alloys do not have adequate corrosion resistance for applications above ambient temperature. The literature does not give any guidance as to the existence of a threshold temperature, above which the material will corrode excessively. In summary, magnesium alloys have their poorest corrosion behavior in immersed, aqueous environments. Generally, the corrosion rates are reported in terms of weight loss rather than penetration depth, and this can be misleading because this material is prone to pitting rather than uniform attack. As a caveat, the corrosion behavior of magnesium alloys under any immersion condition should be regarded with suspicion until proven safe. D. Localized and General Corrosive Attack (Gas Phase and Partial Immersion) Under most conditions, a partial immersion (part gas phase and part liquid phase) is at least as damaging as a full immersion, so that all the observations applicable to full immersion are relevant to a partial immersion. There is the additional complication that the current lines are focused at the liquid side of the gas–liquid phase, resulting in a knifeline attack that may increase the penetration rate. The presence or absence of oxygen does not appear to affect the process. Magnesium alloys have insignificant corrosion rates in clean air if the relative humidity is below about 65% (2); if the surface remains clean of salt deposits, the alloys can remain corrosion-free up to relative humidities of ⬃90%. When the surface becomes contaminated with dirt (which contains salts), the corrosion rates increase and the damage is revealed as shallow pits. Benchmark results for a dirty industrial environment are shown in Table 5. These results can be quite variable, with a rural environment having lowest corrosion rates and a marine environment might be somewhat greater. These results indicate that air exposures are relatively benign for magnesium alloys compared to total immersion. Dry gases such as chlorine, iodine, bromine, and fluorine result in little corrosion of magnesium alloys under ambient temperature conditions. However, the addition of moisture can greatly increase the corrosion rate. Magnesium alloys have linear reaction kinetics in oxygen gas at elevated temperatures and these rates can be very high. For instance, as a benchmark rate, Leontis and Rhines (24) have measured a penetration rate of about 1.4 mm/year for a 9% Al–magnesium alloy at 400°C. The rate would be lower in air and at lower temperatures, but it is not clear what would constitute safe operating conditions; this would have to be determined on a case-by-case situation. Magnesium alloys can be used successfully under inert or low-reactivity environments. There is a successful application (350°C) in the gas-cooled nuclear industry for 266 Danielson Table 5 Penetration Rates for Various Magnesium Alloys Exposed for 2 Years in an Industrial Atmosphere Alloy A8 AZ91 Z5Z ZRE1 TZ6 A8, high purity Penetration rate (mm/year) 0.049 0.047 0.069 0.074 0.084 0.058 Source: Ref. 2. zirconium- or beryllium-containing magnesium alloys where they are used to encapsulate uranium fuel. Clearly, this has to be under inert gas conditions so that the penetration rate by the oxidation process is essentially zero (2). Magnesium alloys are reported to be very corrosion resistant in elevated-temperature H 2 S or SO 2 environments (25). Oxidation rates are low in BF 3 and SF 6 , and these gases are used as cover gases in foundry operations (22). Usually, magnesium alloys are not used at these elevated temperatures because of a significant loss of creep strength. In summary, magnesium alloys have excellent corrosion behavior in lowhumidity environments and at elevated temperatures under oxygen-free atmospheres. Partial immersion and high-humidity environments must be examined on a case-by-case basis. E. Stress-Corrosion Cracking and Corrosion Fatigue (Aqueous Solutions) A serious limitation to the structural use of magnesium alloys is their propensity for SCC (stress-corrosion cracking) in a wide variety of environments, alloy compositions, and metallurgical conditions. Most of the cracks are transgranular and initiate at a surface pit or defect. The SCC review by Miller et al. (26) remains the best overview in this field to the year 1991, and it will be quoted extensively in this section. The Al-containing alloys are considered the most susceptible to SCC, and this susceptibility increases with Al content. Zinc also increases the susceptibility, and the most commonly used class of alloys, AZ, is considered to be particularly susceptible. Those alloys containing no aluminum or zinc are the most resistant, with M1 (Mn–magnesium alloy) being one of the most resistant. SCC is observed in almost any aqueous environment, including deionized water. Magnesium Alloys 267 Halide ions are particularly aggressive toward SCC, but almost any dissolved salt increases the rate over that observed in deionized water. The environments that do not induce SCC are those in which the magnesium alloys are perfectly passive, such as dilute alkalies, concentrated HF, and chromic acid. Nitrates and carbonates also have an inhibiting effect. The threshold concentrations at which dissolved ions can cause SCC are not well defined in the literature, but dilute solutions are capable of accelerating SCC. A popular laboratory test environment for ranking alloys is one containing NaCl ⫹ K 2Cr 2O 4 because it results in rapid SCC. Unfortunately, the results obtained in this environment have not correlated well with field failure experience. The literature is filled with conflicting observations about the effects of wrought versus cast alloys, the effects of cold work, and the effects of heat treatment on susceptibility to SCC. Curiously, the cracking process is slowed or even stopped by cathodic polarization while it is increased with anodic polarization (relative to the open-circuit electrochemical potential); yet, the cracking process is hypothesized to be driven by hydrogen embrittlement. Fracture-mechanics-based SCC crack growth rate data are uncommon in the literature, but some benchmark fracture mechanics data for this process are shown in Fig. 6 in a work by Speidel et al. (27) with a somewhat SCC-resistant alloy, ZK-60A [Al-free alloy containing zirconium, yield strength ⫽ 296 MPa (43 ksi)]. The cracking rate in the two salt solutions was greater than 3 ⫻ 10 ⫺4 cm/s, a very high rate, with the K 1scc ⱕ 7.7 MPa m 1/2 (7 ksi in.1/2 ). The effects of the dissolved salt are to increase the cracking rate and lower the K 1scc relative to the distilled-water environment. The authors make the important comment that in their experience, all high-strength magnesium alloys behave similarly to environment-enhanced subcritical crack growth and that this behavior does not vary significantly with alloy composition and heat treatment. Both SCC and corrosion fatigue cracks propagated in a mixed transgranular and intergranular mode in this testing. The same environments that cause SCC also reduce the corrosion fatigue (CF) performance. Magnesium alloys have their best CF performance in vacuum, followed by air, deionized water, and salt solutions. The CF strengths can be as low as 10% of those in air. As with SCC, the aluminum- and zinc-containing alloys are especially susceptible, even in mildly corrosive environments (28). A significant amount of S–N data exist in the literature, but most of it is old. The majority of the modern fracture-mechanics-based fatigue crack growth data comes from Russia. The effects of wrought versus cast alloys, the effects of cold work, and the effects of heat treatment are largely unexplored with CF. Some fracture-mechanics-based (da/dN–∆K) benchmark data for CF are shown in Fig. 7 (27). This work clearly demonstrates that salt solutions significantly increase the CF rate and that distilled water is more aggressive than dry argon (particularly in stage III crack growth). In a more recent study, Stephens et al. (29) examined the CF response of a modern, high-purity sand-cast alloy, AZ91E-T6, in a 3.5 268 Danielson Fig. 6 The effect of stress intensity and environment on the velocity of SCC cracks for a high-strength magnesium alloy. wt% NaCl solution and air (R ⫽ 0.05 and 0.5, respectively). Their CF testing was limited to the stage II regime, but the results were very similar to the results of Spiedel et al., implying that their earlier comment about the similarity of all magnesium alloys with regard to SCC behavior might also be extended to their CF behavior. The threshold ∆K for CF is ⬍2 MPa m 1/2 which is even lower than the K1scc for SCC. These are very low values. Increasing the temperature generally acts to increase the propensity for SCC and CF, but the literature data are extremely sparse on this variable (6). In summary, the SCC and CF behavior of magnesium alloys is seriously Magnesium Alloys 269 Fig. 7 The effect of cyclic stress intensity range and environment on the velocity of CF cracks for a high-strength magnesium alloy. degraded by immersion in aqueous solutions and probably imposes a greater limitation to the structural use of these alloys than their general and localized attack properties. Careful mechanical design (staying below the K 1scc or threshold ∆K) will be needed to successfully use magnesium alloys for structural components. F. Stress-Corrosion Cracking and Corrosion Fatigue (Gas Phase and Partial Immersion) Magnesium alloys have their best CF behavior in a vacuum environment. The CF strength is lower in dry air and continues to decrease as the relative humidity 270 Danielson increases. Bothwell (16) reports that an AZ31 alloy lost CF strength once the relative humidity exceeded 50%, and at 93% relative humidity, it had a fatigue strength of about 75% that of dry air. G. New Processing Methods Two rapid-cooling methods have been utilized to improve the corrosion resistance of magnesium alloys: rapid solidification (RS) and laser annealing (LA). Both methods are based on rapid cooling of the melt such that metastable, supersaturated solid solutions are obtained. The basic idea is to minimize the amount of the various intermetallic compounds that could act as initiation sites for localized corrosion processes. The RS methods result in the formation of fine powders or foils that must be later extruded into structural forms. LA takes place after the part is nearly machined to the final dimensions, whereupon a high-energy laser beam is focused on the surface and melted to a depth of a few micrometers, and then the part is machined to the final dimensions. The work of Chang and coworkers (30,31) demonstrated the improved general and localized corrosion resistance of the RS alloys as compared to conventional alloys in 3% NaCl. These studies were carried out on RS-formed Mg–Al–Zn alloys with Mn, Si, and rare earth additions. The materials were extruded into a bar before corrosion testing. The measured corrosion rate for the Mg–Al–Zn–Y alloy was about 400 µm/ year (10 mpy), one of the lowest rates ever measured for a magnesium alloy in a salt solution. Makar et al. (32) studied the SCC of RS Mg–Al alloys in NaCl– K 2 CrO4 solution using the slow-strain-rate method. The materials showed the classic transgranular SCC morphology, but no crack growth rate data were reported, so that a comparison could be made with conventional materials. The RS alloy showed an improved general corrosion behavior and very rapid repassivation kinetics when the electrode was scratched. The LA technique has been applied to the commercial alloy AZ91C by first depositing a thin layer of Al, Cr, Zn, and so forth on the surface, followed by laser annealing to melt the surface in order to incorporate the coating into the surface alloy layer (33,34). Both Zn and Al deposits resulted in an improved general corrosion rate behavior, but this was determined by potentiodynamic scans which measures only short-term behavior. Once the thin modified layer is perforated by corrosive attack, the material will probably revert to its bulk behavior. Another new surface modification technique is ion implantation (5). With this technique, the part must be placed in a vacuum chamber and ions (to be implanted) are accelerated and ‘‘slammed’’ into the surface to a penetration depth of 50–500 nm. The method is only mentioned here for information because the penetration depth is extremely shallow and the results have had little effect on the corrosion rate. Magnesium Alloys 271 Magnetron sputter deposition is another surface-modification technique for two or more alloying components. Miller et al. (26) have fabricated nonequilibrium Y–Mg alloys in the 9–22 at% Y range in which the Y remained in solid solution. The metal films were very thin (⬃2 µm thick). Potentiodynamic scans with a potentiostat in 0.1M NaCl indicated improved corrosion resistance over that of conventional magnesium alloys. In summary, both the RS and LA alloys have improved the general corrosion behavior, but there is a lack of SCC and CF rate data that would demonstrate any superiority over conventional materials. Coatings can increase the time for failure from localized attack and potentially for SCC and CF processes by increasing the initiation time of the degradation process, but once the coating is breached, the propagation phase will likely proceed at the same rate as in the conventionally bulk material. H. Coatings Coatings are useful in improving the cosmetic appearance of finished products and extending the service life under certain marginal corrosive environments. There are a great many coating schemes in practice which fall into two broad categories although they are complementary: chemical or electrochemical methods which produce a corrosion-resistant film and coatings which require the application of an impervious top coat. The reader is directed to Refs. 2 and 22 for a comprehensive coverage of this topic. Briefly, the coating process involves cleaning the magnesium-alloy surface, first by degreasing followed by chemical or electrochemical etching (anodizing). Etching often has the beneficial effect of removing the surface-exposed intermetallic particles or heavy metal contamination (i.e., embedded by grit blasting) which can act as cathodes and drive the galvanic attack. Next, some type of conversion coating is applied which may contain a corrosion inhibitor (i.e., chromium), which prepares the surface for adhesion of an organic top coat (i.e., paint). The top coat serves as a barrier between the metal and the outside environment. The creation of the corrosionresistant film follows a similar processing scheme (although certain details such as the composition of the anodizing solution may differ), except that there is no top coat applied. Once the coating is breached by wear or abrasion, the bare metal will be subject to the corrosion processes at the uncoated rate. Coatings would be very useful in certain marginal environments such as alternate wet/dry exposure where general and localized attack are problems but would be dangerous in environments where the bare alloy has a serious corrosion problem, unless extra effort is directed toward frequent inspection and repair of the coating. Coatings for alloys used in structural applications may give inadequate protection from SCC and CF degradation processes because they only extend the initiation time and could act to impede the visual inspection that would be necessary. 272 I. Danielson Macroscopic Galvanic Attack The presence of intermetallic compounds, particularly iron compounds, results in an extremely important form of microgalvanic attack that has been examined extensively in this review. A macroscopic form of this galvanic attack is similarly present when a Mg alloy is fastened (coupled) to a metal of different composition. Although increasing the purity of the alloys has had a very beneficial effect on the microgalvanic attack processes, it is of no avail for the macrogalvanic process. This is a very serious, practical problem which requires attention to detail for the immersed or occasionally immersed condition. Froats et al. (22) have written about this subject in a comprehensive manner, and Olsen (35) has focused on galvanic corrosion of automobile driveline components to which the reader is directed. Basically, prevention of macrogalvanic attack requires attention to the following: (1) minimizing the entrapment of water, (2) selecting metals in direct contact with the Mg alloys that have high hydrogen overvoltages, (3) inserting electrical insulators between the metallic fasteners, if possible, and (4) using protective coatings. Galvanic attack is least with pure aluminum (aluminum alloys are somewhat worse) in electrical contact with a Mg alloy, and carbon and stainless steels are the worst. A compromise solution might be to use an aluminumalloy washer between the steel fastener and the Mg–alloy component. Applying a coating over the cathode (i.e., steel fastener) helps to reduce the surface area; hence, the electrochemical driving force for the galvanic process. By attention to detail, it is possible to make acceptable compromises that result in adequate performance of Mg alloys. The effect of dissimilar materials is unknown on SCC and CF processes. IV. CONCLUSIONS Magnesium alloys have very desirable structural properties from the standpoint of weight and strength, but these properties are offset by their unusual propensity for corrosive attack by general, localized, stress-corrosion cracking, and corrosion fatigue under aqueous environments. The first line of defense is using high-purity alloys containing low amounts of Fe, Cu, and Ni. There appears to be no inhibitor for these degradation processes in most engineering environments. Coatings can improve the cosmetic appearance and extend the useful operation life in marginal environments. Accelerated testing in the laboratory has a history of not predicting field experience, and the technical literature usually reports test results in these accelerated environments. Clearly, laboratory testing must realistically replicate the application environment, particularly with respect to SCC and CF processes, if these alloys are to be economically and safely utilized. Magnesium Alloys 273 REFERENCES 1. MO Pekguleryuz, MM Avedesian. Magnesium alloying—Some metallurgical aspects. In: BL Mordike, F Hehmann, eds. Magnesium Alloys and Their Applications. Oberursel: Informationgesellschaft-Verlag, 1992, pp. 213–220. 2. KG Adamson, DS Tawil. Magnesium and magnesium alloys. In: LL Shreir, RA Jarman, GT Burstein, eds. Corrosion, 3rd ed. Boston: Butterworth/Heinemann, 1994, pp. 4:98–4:115. 3. JE Hillis. Magnesium. In: R Baboian, ed. Corrosion Tests and Standards: Application and Interpretation. Philadelphia: ASTM, 1995, pp. 438–446. 4. A Luo, MO Pekguleryuz. Cast magnesium alloys for elevated temperature applications. J Mater Sci 29:5259–5271, 1994. 5. GL Makar, J Kruger. Corrosion of magnesium. Int Mater Rev 38:138–153, 1993. 6. WK Miller. 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INTRODUCTION Many investigations have shown that low ductility and brittle fracture in intermetallics are caused not only by intrinsic factors (such as lack of sufficient deformation modes, poor cleavage strength, weak grain boundaries, etc.) but also by extrinsic factors. Environmental degradation, an extrinsic factor, is found to be a major cause of brittle fracture in many ordered intermetallics, particularly those with high crystal symmetries (i.e., cubic L12 and B2 and hexagonal DO19), as outlined in several recent reviews (1–3). There are two types of environmental embrittlement observed in intermetallics. One is hydrogen-induced embrittlement occurring at ambient temperatures. Many intermetallic alloys show a substantial decrease in room-temperature tensile ductility due to moisture-induced hydrogen embrittlement in air, or as a result of direct exposure to hydrogen by precharging or testing in hydrogen gas. The other is oxygen-induced embrittlement in air at elevated temperatures. In most cases, the embrittlement is due to a dynamic effect involving generation and penetration of an embrittling agent (i.e., hydrogen or oxygen) during testing. In this chapter, the two types of environment-sensitive degradation are treated separately. The concluding section discusses metallurgical means of alleviating environmental degradation. 275 276 Stoloff II. EMBRITTLEMENT AT AMBIENT TEMPERATURES Hydrogen, introduced by exposure to moist environments, by cathodic charging, or by testing in hydrogen gas, has been shown to embrittle a large number of single-phase intermetallics. Table 1 lists those nickel- and iron-containing intermetallic alloys that have been tested for susceptibility to environmental embrittlement (2). Most notable is the susceptibility to moisture or hydrogen of all undoped L12 intermetallics that have been examined to date, except for Ni3Fe, which is embrittled only when hydrogen is charged into the alloy. Therefore, it is highly likely that all L12 intermetallics containing substantial amounts of the transition metals Fe, Ni, or Co will prove to be susceptible to contact with hydrogen or hydrogen-containing environments. Except for the iron aluminides, data for other intermetallics remain sparse. A. L12 Alloys In terms of environmental embrittlement, ordered L12 intermetallics can be grouped into two categories: (a) alloys containing no reactive elements and (b) alloys containing reactive elements (e.g., Al or Si). In the first group, the alloys are severely embrittled only when charged with hydrogen, such as by cathodically charging. For the second category, the alloys themselves are capable of generat- Table 1 Alloys Embrittled by Moisture or Hydrogen at Ambient Temperature Alloy Ni3(Al,Ti) (single crystal) Ni3Al⫹B Ni3Al⫹Be Ni3(Al,Mn) Ni3Si Ni3(Si,Ti) Ni3(Si,Ti)⫹B (Co,Fe)3V Ni3Fe FeAl Fe3Al Crystal structure Environmental embrittlementa Ref. L12 䊊 4 L12 L12 L12 L12 L12 L12 L12 L12 B2 DO3 䊊 䊊 䊊 䊊 䊊 ⫻ 䊊 ⫻ 䊊 䊊 5 6 7 8 6 9 10 11 12 䊊: Environmental embrittlement observed in moist environments; ⫻: not observed in moist environments, but embrittled by hydrogen. Source: Ref. 2. a Embrittlement of Ni- and Fe-Based Intermetallics 277 ing hydrogen from hydrogen-containing environments at ambient temperatures. The most striking case is severe embrittlement of iron aluminides in moist air at room temperature. 1. Ni3Fe Ni3Fe is a model intermetallic for the study of hydrogen embrittlement in alloys displaying a critical ordering temperature, Tc. This has permitted a direct comparison to be made between the ordered and disordered conditions of Ni3Fe, as shown in Fig. 1 (13). In each case, hydrogen was cathodically charged for 1 h prior to tensile test at room temperature. Ni3Fe is slightly embrittled by precharging in both the ordered and disordered conditions; simultaneous charging and testing causes greater loss of ductility than precharging, with an especially severe effect for the ordered condition, perhaps due to the transport of hydrogen by moving dislocations. Fractographic studies showed a change in fracture mode from transgranular microvoid coalescence to intergranular fracture in hydrogen accompanying the reduction in elongation. However, the intergranular embrittlement zone was only about one-third as deep in precharged samples of ordered Ni3Fe as in disordered samples. Chia and Chung (14) report that Ni3Fe is not embrittled by moisture in air or by water, in marked contrast to the effects of hydrogen. 2. Ni3 Al a. Effects of Composition Although the embrittling effects of hydrogen on fracture behavior of Ni3Al alloys have been well documented (6,15), only recently has evidence been provided for moisture-induced cracking in these alloys. Liu (16) has reported that the tensile ductility of Ni-24 at% Al is increased from 2.6% to 7.2% by testing in oxygen rather than in air. A similar result was noted with Ni–23.5 at% Al (2.5% versus 8.2% elongation in air and oxygen, respectively). Higher ductilities are observed in vacuum than in air for boronfree Ni–23.4 at% Al, but the difference is strain rate sensitive; see Fig. 2 (17). Also, the vacuum level plays a role, with the ductility of polycrystalline Ni–23.4 at% Al increasing from 7.9% to 23.4% as the vacuum level changes from 10⫺1 to 3.6 ⫻ 10⫺8 Pa (17). Additional striking effects of the environment have been reported for a Ni– 22.65 Al–0.26 Zr alloy, recrystallized from a 〈110〉-oriented single crystal and tested at room temperature in two orientations; see Table 2 (18). This boron-free alloy displayed 8.7% elongation in water, 13.2% in air, and 50.6% in oxygen in the 0° orientation. Fracture was predominantly intergranular in all environments. Although the results could have been influenced by texture (unknown in these alloys) or by a preponderance of fracture-resistant grain boundaries in this material, the authors suggested that the principal role of boron in ductilizing Ni3Al 278 Fig. 1 Stoloff Embrittlement of Ni3Fe by cathodically charged hydrogen. (From Ref. 13.) Embrittlement of Ni- and Fe-Based Intermetallics 279 Fig. 2 Effect of strain rate on room-temperature tensile ductility of polycrystalline Bfree Ni3Al (23.4% Al) in air and vacuum. (From Ref. 17.) may be to suppress environmental embrittlement. Grain-boundary strengthening, permitting transgranular failure, would be a secondary effect. [Another factor is the zirconium content, as it has been shown that zirconium improves the ductility of Ni3Al in air and, especially, in oxygen (19).] Additional evidence that boron suppresses environmental embrittlement is summarized in Table 3, which shows that undoped Ni–23.4% Al is severely embrittled in air, whereas a Ni–24 Al–0.2 B alloy displays high ductility in air, water, and oxygen (17). These results are consistent with an earlier report by Takasugi et al. (5), who showed that the ductility of boron-doped Ni3Al was nearly the same in vacuum and in air. The minimum amount of boron required to eliminate environmental embrittlement is between 50 and 100 wppm. Wan et al. (20) have reported that the tensile ductility of Ni–23 at% Al– 120 wppm B is only 0.6% in 0.1 MPa hydrogen gas and rises to 10.1% in air. However, when tested in 1.3 ⫻ 10⫺3 Pa vacuum or in 0.1 MPa oxygen, ductilities of 20% or more are achieved. Although preoxidation can suppress embrittlement by cathodically charged hydrogen (15), embrittlement by hydrogen gas was little affected by preoxidizing for 24 h at 900°C (20). Previous work by Takasugi and Izumi (8), carried out on Ni3Al with 500 wppm B, had shown no effect of the test environment on ductility. Wan et al. (20) concluded that higher boron contents may suppress grain-boundary decohesion in the presence of hydrogen released from water vapor. Clearly, however, boron does not prevent hydrogen embrittlement under cathodic charging conditions (15,21). 280 Table 2 Effects of Environment and Specimen Orientation on Tensile Properties of Ni–22.65% Al–0.26% Zr, 25°C, ε ⫽ 5.3 ⫻ 10⫺3 s⫺1 Specimen orientation 0° 45° Test environment Elongation to fracture (%) Yield strength (MPa) Ultimate tensile strength (MPa) σ10 ⫺ σ5a (MPa) σ15 ⫺ σ5a (MPa) Water Air Oxygen Water Air Oxygen 8.7 13.2 50.6 6.3 10.7 47.8 322 324 326 331 341 327 528 661 1451 473 603 1438 — 133 127 — 127 143 — — 250 — — 271 a Difference in flow stress at 10% and 5% (or 15% and 5%) strain. Source: Ref. 18. Stoloff Effect of Test Environment on the Room-Temperature Tensile Properties of B-Free and B-Doped Ni3Al Alloy chemistry (at.%) Ni–23.4 Al Ni–24 Al–0.2 B Test environment Elongation to fracture (%) Yield strength (MPa) Ultimate strength (MPa) Air Oxygen UHV Water Air Oxygen 3.1 15.8 23.4 36.8 39.3 42.8 308 336 —a 288 280 289 392 681 — 1199 1241 1315 Embrittlement of Ni- and Fe-Based Intermetallics Table 3 a Not measured. Source: Ref. 17. 281 282 Stoloff This observation suggests that hydrogen, generated either by the moisture– aluminum reaction or by hydrogen charging, diffuses mainly along grain boundaries and causes intergranular fracture. Note that a similar reduction in ductility and change in fracture mode has been observed in the other L12 intermetallic alloys listed in Table 1. The role of boron in suppressing environmental embrittlement of Ni3Al alloys has been studied extensively, but the precise mechanism of enhancement of environmental resistance remains elusive. Cohron et al. (22,23) have shown that at three different boron levels, 50, 150, and 500 wppm, the ductility of Ni3Al decreases with increasing pressure of hydrogen gas; see Fig. 3. The 500 wppm boron alloy had previously been considered to be immune to the environment. A shift in fracture mode from transgranular to intergranular is observed as the ductility drops, suggesting that the level of intergranular fracture correlates well with ductility in this intermetallic. The same study showed that a boron-free alloy was more ductile than alloys containing 50 or 100 wppm boron at hydrogen Fig. 3 Elongation to fracture of three B-doped Ni3Al (24% Al) alloys as a function of hydrogen pressure. Open and closed symbols refer to the ion gauge on and off, respectively. (From Refs. 22 and 23.) Embrittlement of Ni- and Fe-Based Intermetallics 283 pressures exceeding about 1 Pa. Because atomic hydrogen is the cause of embrittlement, this surprising result suggests that boron promotes the dissociation of molecular to atomic hydrogen. This study also showed that when all sources of hydrogen were minimized, the ductility of polycrystalline Ni3Al exceeded 40%, with a predominantly transgranular fracture path. Zirconium, as cited previously, is another important dopant that can reduce environmental embrittlement in Ni3Al. George et al. (19,24) and Lin et al. (25) suggested that zirconium enhances grain-boundary cohesion. Chiba et al. (26), on the other hand, proposed that zirconium reduces the extent of sulfur segregation to grain boundaries, thereby inhibiting sulfur-induced loss of cohesion. Itoh et al. (27) recently attempted to resolve these conflicting views by studying the ductility of sulfur-doped Ni3Al containing different levels of zirconium. Although they confirmed that zirconium alleviated environmental embrittlement and found that ZrS particles formed when sulfur levels were high, they could not definitively identify the mechanism of the improved behavior. Zirconium has a similar beneficial effect on ductility and fatigue crack growth rates in Fe3Al alloys [see Sec. II.C.2, (28)], but it is difficult to see how environmental effects in the two alloy systems are related because moisture induces intergranular fracture in Ni3Al but does not change the characteristic transgranular crack path in Fe3Al. Clearly, much more work has to be done to identify the cause of the benefits bestowed by small additions of zirconium. Apart from boron and zirconium, several other solutes can increase the ductility and change the crack path of polycrystalline Ni3Al in air; see Table 4 (29). Transition metals such as chromium, iron, and manganese improve ductility Table 4 Effect of Alloying Additions on Room-Temperature Ductility and Fracture Behavior of Ni3Al Alloys Prepared by Conventional Melting and Casting Alloy element B B,Fe Mn Fe Pd Pt Co Cu Zr Source: Ref. 29. Alloy composition (at.%) Tensile ductile (%) Fracture mode Ni3Al (24% Al) Ni–24 Al–0.5 B Ni–20 Al–10 Fe–0.2 B Ni–16 Al–9 Mn Ni–10 Al–15 Fe Ni–23 Al–2 Pd Ni–23 Al–2 Pt Ni–23 Al–2 Co Ni–23 Al–2 Cu Ni–22.65 Al–0.26 Zr 1–3 35–54 50 16 8 11 5 4 6 13 Intergranular Transgranular Transgranular Transgranular Mixed Intergranular Intergranular Intergranular Intergranular Intergranular 284 Stoloff in air, but the effects of these elements may be due to the introduction of the disordered γ-phase in the microstructure. Recently, palladium has been reported to increase ductility in air of single-phase Ni3Al (30). Palladium in the range 1– 3 wt% also increases ductility in the presence of hydrogen, although considerable scatter in the data was noted (31). Further, Pd-free Ni3Al containing 5 wppm hydrogen exhibited cleavagelike facets, whereas an alloy containing 0.5% Pd exhibited microvoids on the fracture surface. A similar beneficial effect of Pd has been noted in steels, for which the fracture toughness and threshold stress in constant-load tests were significantly improved by Pd additions (32). The partial replacement of aluminum by titanium in Ni3Al single crystals does not change the pattern of embrittlement by moisture, water vapor, and hydrogen; see Fig. 4 (33). A strain rates approaching 10⫺3, embrittlement disappeared in all environments. This work also demonstrated that when grain boundaries are absent, embrittlement still can occur, along low-index cleavage planes. b. Effects of Predeformation Moisture-induced embrittlement of Ni3Al (34) and Ni3Si,Ti (35) can be reduced by predeformation (e.g., by shot peening or by tensile prestrain), as shown in Fig. 5. The effect varies with temperature of the predeformation; elongation increases with decreasing predeformation temperature. Corresponding to increased tensile ductility, the fracture mode changes from intergranular to transgranular. The effect of prior deformation was attributed Fig. 4 Effects of environment and strain rate on the ductility of 〈100〉-oriented Ni3Al, Ti single crystals. (From Ref. 33.) Embrittlement of Ni- and Fe-Based Intermetallics (a) 285 (b) Fig. 5 Effect of prestrain on environmental embrittlement in Ni3(Si,Ti) polycrystals at room temperature: (a) simply deformed and (b) prestrained up to 11% in liquid nitrogen and then strained to fracture. (From Ref. 35.) to trapping of hydrogen atoms at vacancies or dislocations or an increase of soluble hydrogen in the lattice. c. Stress-Corrosion Cracking Ricker et al. (36) have shown that stresscorrosion cracking of Ni3Al alloyws such as IC-50 (Ni–22 at% Al, 0.01 Zr, 0.08 B) and IC-218 (Ni–17 at% Al, 8.18 Cr, 0.127 Zr, 0.098 B) in acidic solutions (HNO3, H2SO4) arises from the liberation of hydrogen. Slow strain-rate tests in low-pH solutions show a large reduction in ductility and a change in fracture mode from transgranular to intergranular compared to tests in neutral or alkaline solutions. Preexposure to sulfuric acid also reduces ductility in a subsequent slowstrain-rate test conducted in air, again due to release of hydrogen at the specimen surface (36,37). 3. Ni3Si Intergranular fracture and environmental embrittlement have been studied in alloys based on Ni3Si (7,38). Ni3Si showed no appreciable plastic deformation when tested in moist air but displayed an elongation of 7.5% when tested in dry oxygen (Table 5), demonstrating that Ni3Si is prone to environmental embrittlement (39). Because the elimination of the environmental effect by testing in dry oxygen does not lead to extensive ductility (e.g., 30% or more) and complete suppression of intergranular fracture in Ni3Si, moisture-induced hydrogen embrittlement does not appear to be the sole source of grain-boundary brittleness in the silicide. Boron additions segregate strongly to grain boundaries in Ni3Si and suppress environmental embrittlement, as is clearly shown in Table 5 (39). Note 286 Stoloff Table 5 Effect of Test Environment on Room-Temperature Tensile Properties of Ni3Si (22.5% Si) With and Without Boron Additions Test environment Tensile ductility (%) Yield strength [MPa (ksi)] Ultimate tensile strength [MPa (ksi)] 0 ppm Ba Air Vacuum Oxygen b 0 4.7 7.5 677 (98.3) 685 (99.4) 627 (91.0) 853 (124) 1040 (151) 774 (112) 813 (118) 1151 (167) 1268 (184) 655 (95.0) 659 (95.6) 703 (102) 819 (119) 823 (120) 837 (136) 610 (88.6) 590 (85.6) 875 (127) 854 (124) 50 ppm Bc Air Vacuum 8.1 8.9 100 ppm Ba Air Vacuum Oxygen 5.0 5.0 5.9 150 ppm Bc Air Oxygen 7.0 6.6 a Fractured prior to macroscopic yielding. Specimens were annealed 3 days/950°C ⫹ 1 day/600°C. c Specimens were annealed 1 day/950°C ⫹ 1 day/600°C. Source: Ref. 39. b that as little as 50 ppm boron provides essentially the same ductility in air and in vacuum, whereas the undoped alloys display no ductility in air and 4.7% in vacuum. Other elements that reduce the effects of moisture in air or of distilled water on the ductility of Ni3(Si,Ti) alloys are chromium, manganese, and iron, in that order of effectiveness (40). Accompanying improved ductility is a reduction in the extent of intergranular fracture. Reduced strain rates are detrimental to embrittled alloys, as is common with other intermetallics. However, the mechanism of improvement in ductility with transition metal solutes is not clear, with several hypotheses being offered. Furthermore, the addition of these transition elements was less effective in reducing embrittlement induced by hydrogen gas. This implies that decomposition of H2 into H at the surface is not markedly affected by these solutes. B. NiAl Little is known about possible environmental embrittlement of NiAl. Lahrman et al. (41) reported that the tensile properties of NiAl single crystals are about Embrittlement of Ni- and Fe-Based Intermetallics 287 the same in air, in vacuum, and in oxygen at room temperature. However, Bergmann and Vehoff (42) have reported that gaseous hydrogen, charged into single crystals at 1000°C, promotes stable crack growth in 〈110〉-oriented single crystals at temperatures above 475°K. Extensive scatter in the data precluded any conclusions as to the possible influence of hydrogen on the toughness of polycrystals. C. Iron Aluminides The two iron aluminides, FeAl and Fe3Al, form back-centered-cubic (bcc)-ordered crystal structures of B2 and DO3 type, respectively. Despite the difference in crystal structure, these aluminides show somewhat similar patterns of embrittlement in hydrogen-containing atmospheres at room temperature. Both aluminides exhibit only a few percent ductility (1–5%) when tested at ambient temperature in air. This led many researchers to conclude that these alloys are inherently brittle. However, it has been demonstrated conclusively that when water vapor and hydrogen are eliminated from the external environment, both alloys exhibit considerable ductility. 1. FeAl Tensile elongations up to 19% have been achieved in FeAl alloys tested in ‘‘dry’’ environments such as oxygen (11). An increase in ductility from 2% in water to 18% in oxygen was accompanied by a change in fracture mode from transgranular cleavage in air to mainly grain-boundary separation in dry oxygen. These observations suggest that cleavage planes are more susceptible to embrittlement than are grain boundaries. The maximum degree of moisture-induced embrittlement occurs on either side of room temperature; see Fig. 6 (1). At higher temperatures, in situ protective oxide films can form readily on specimen surfaces, whereas at lower temperatures, the aluminum–moisture reaction is slowed and the equilibrium moisture content in air also is lowered. a. Effects of Aluminum Content The ductility of FeAl decreases with increasing aluminum content in the range 38–48 at%, as shown in Fig. 7 (43). Environmental embrittlement of FeAl, defined as the difference in ductility between tests in air and in oxygen, is considerably reduced as the aluminum content increases. For Fe–40% Al, ductility is about 4% in air, in 2 ⫻ 10⫺7 torr vacuum, or after hydrogen charging (44). Gaydosh and Nathal (45) reported that the ductility of Fe–40% Al is sensitive to microstructure (annealed and furnace-cooled material displayed 9% elongation in vacuum compared to 5% for an as-extruded sample). However, boron increased ductility of as-extruded Fe–40% Al to 9% in vacuum. Fe–50% Al, on the other hand, was brittle in vacuum for both the as-extruded and furnace-cooled conditions. For Fe–43 at% Al, the ductility is nil in air as well as in dry oxygen; all specimens fail intergranularly. This differ- 288 Stoloff Fig. 6 Effect of test temperature on tensile elongation of FeAl (36.5% Al) in air. (From Ref. 1.) Fig. 7 Effect of aluminum concentration and test environment on the room temperature tensile ductility of B-doped and B-free FeAl. (From Ref. 43.) Embrittlement of Ni- and Fe-Based Intermetallics 289 ence in behavior with aluminum content suggests that grain boundaries in FeAl alloys with Al ⬎ 38% are intrinsically weak. Therefore, environmental embrittlement and intrinsic effects must be distinguished in order to establish strategies for reducing brittleness. It has been demonstrated that the intrinsic grain-boundary brittleness in FeAl and other intermetallics can be alleviated by microalloying with boron (45,46). Figure 7 shows that environmental embrittlement does not occur in borondoped Fe–48% Al, but ductility is low in both inert and aggressive environments (42). As aluminum content decreases, embrittlement is increasingly severe in boron-doped alloys, as measured by the difference in ductility between oxygen and air. It is clear from Fig. 7 that the intrinsic ductility of boron-doped FeAl drops sharply with increasing aluminum content, paralleling observations on boron-free material cited earlier. b. Effects of Annealing Temperature It has been widely reported that furnace cooling from elevated temperatures or the imposition of low-temperature anneals markedly increases the ductility of FeAl alloys (47–51). This effect has been linked to the elimination of excess vacancies, which cause hardening, as a result of these treatments. The influence of annealing temperature on environmental embrittlement of Fe–36% Al has been studied by Lynch and Heldt (49,50). Ductility in air is much lower than in dry oxygen at all annealing temperatures, but there is a detrimental effect of raising annealing temperature in either environment. It was concluded that excess vacancies resulting from cooling affect the intrinsic ductility of iron aluminides and not their susceptibility to moisture. c. Effects of Grain Boundaries There have been no systematic studies of the influence of grain size on the susceptibility of FeAl alloys to environmental embrittlement. However, smaller grain sizes tend to favor increased ductilities for similar heat treatments, as noted by both Klein and Baker (51) and Gaydosh et al. (52). Lynch et al. (53) have shown that the tensile elongation of single crystals of Fe–35% Al is extremely sensitive to the environment; see Fig. 8. Note that the ductility in air is about 4%, independent of orientation, whereas ductilities in oxygen are 18.5% for crystal No. 1, oriented with the tensile axis 40° from the (100) pole and 22.7% for crystal No. 2, oriented with the tensile axis 40° from the (100) pole. These differences in ductility are similar to those observed for polycrystals, suggesting that grain boundaries are not necessary for moistureinduced embrittlement of FeAl alloys. (The same is true for Fe3Al alloys, in which transgranular cleavage is usually observed in air as well as in oxygen or vacuum.) The similarities between water-vapor-induced embrittlement and hydrogen embrittlement are now well established for both FeAl and Fe3Al alloys. For example, lower strain rates exacerbate embrittlement in air, exactly as would be expected for a hydrogen-related phenomenon (54,55). Recently, Kasul and Heldt 290 Fig. 8 Stoloff Elongation of Fe–35 Al single crystals in oxygen and air. (From Ref. 53.) (56) have provided additional evidence to link both types of embrittlement. Subcritical crack velocities were measured in Fe–35 at% Al under constant-load conditions; these were compared with strain-rate effects on ductility in air, as well as with results of cathodic charging experiments. It was shown that the strain rate has essentially no effect on ductility in 10⫺3 Pa vacuum, whereas ductility was sharply reduced with decreasing strain rate in air (56). Similar results in air have been reported by Shan and Lin (55), but they are at variance with a report by Nagpal and Baker (54) of a sharp drop in ductility at a particular strain rate. At a strain rate of about 0.3 s⫺1 the ductilities in air and vacuum were equal; this strain rate corresponded reasonably well with the established critical strain rate above which enhanced hydrogen transport on moving dislocations is no longer likely (56). The role of moving dislocations in transporting hydrogen to potential crack initiation sites has been discussed by several research groups (13,56), but there is, as yet, no consensus that this mechanism is applicable to environmental embrittlement phenomena. d. Cathodic Charging Experiments Cathodic charging experiments on Fe–35 at% Al have shown a clear correlation among hydrogen content, charging time, and ductility (56). For example, after 24 h of charging in 1N H2 /SO4, the hydrogen content is about 1.3 ppm and the ductility in subsequent tests in vacuum has dropped from 8% to 1.3% (56). Baking treatments of sufficient duration at high temperature (800°C) had previously been shown by Kasul and Heldt (57) Embrittlement of Ni- and Fe-Based Intermetallics 291 to completely restore the ductility of cathodically charged Fe3Al. For Fe–35% Al, baking at 400°C for more than 1 h was sufficient to completely restore ductility. A valuable result of this study was an estimate of the diffusivity of hydrogen in Fe–35% Al at room temperature: about 4 ⫻ 10⫺12 cm2 /s. Because this diffusivity is too low to allow significant hydrogen penetration ahead of a crack moving under constant stress in stage 2, it was suggested that hydrogen transport by moving dislocations may occur. e. Fracture Toughness Relatively few studies of fracture toughness in the iron aluminides have been reported. Klein et al. (58) have shown that the fracture toughness of FeAl increases from 23 to 30 MPa m1/2 for blunt notched specimens as the strain rate increased. A similar trend was noted for precracked samples. Ko et al. (59) used the standard multispecimen JIc procedure on side-grooved specimens to estimate fracture toughness of Fe–35 at% Al as a function of test environment. It was shown that toughness increased from air to vacuum to oxygen environments; see Fig. 9. Also shown are toughness data for an Fe3Al alloy, Fe–28 at% Al, tested in both the B2 and DO3 conditions. For both conditions, toughness in the vacuum and oxygen environments was lower than for the FeAl Fig. 9 Fracture toughness of iron aluminides in different environments. (From Ref. 59.) 292 Stoloff alloy, whereas tests in air resulted in the lowest toughness in FeAl. These results were very similar to tensile data reported for the two alloys by the same group. Under both tensile and fracture toughness loading conditions transgranular cleavage was the dominant fracture mode seen in air. Some intergranular fracture regions were seen in B2-ordered Fe–28 Al and Fe–35 A1 specimens tested in oxygen. This study indicated that Fe–35 Al is more susceptible to environmental embrittlement than is Fe–28 Al and that grain boundaries are intrinsically weaker in the FeAl alloy. However, complicating the interpretation of the data is the fact that the B2 condition of Fe3Al alloys is more ductile and displays higher toughness than the DO3 condition. This suggests that it is the higher Al content in the Fe–35 Al alloy, rather than the crystal structure, that is responsible for the greater effect of the environment. f. Fatigue Crack Growth Recent work in our laboratory has shown that hydrogen severely embrittles FeAl under cyclic loading conditions; see Fig. 10 (60). It appears that moisture is embrittling only at high cyclic stress intensities for this alloy. 2. Fe3 Al Fe3Al alloys form in the range of about 18–32 at% Al. When near-stoichiometric alloys are quenched from above the critical ordering temperature, Tc, they display a partially ordered B2 structure. Slow cooling through Tc or extended annealing Fig. 10 Effects of environment on fatigue crack growth of Fe–35 at% Al at 25°C. (From Ref. 60.) Embrittlement of Ni- and Fe-Based Intermetallics 293 just below Tc results in a highly ordered DO3 structure at room temperature. There is considerable evidence that environmental embrittlement, resulting from moisture in air, is severe in binary Fe3Al alloys. For example, tensile elongation in air is 4% but 19% in vacuum, whereas less than 1% strain is noted after precharging with hydrogen, as shown in Fig. 11 (44). Ductility of Fe–28% Al alloys in air is significantly increased by the addition of chromium (61). Although the mechanism of improvement in ductility is unknown, it may result from modification of the Al2O3 coating that naturally forms on Fe3Al in such a way that hydrogen liberation from water vapor is reduced. It is unlikely that Cr2O3 replaces Al2O3 as the protective oxide at these chromium levels. Unfortunately, chromium additions do not suppress embrittlement by gaseous hydrogen or by hydrogen introduced by electrolytic charging (44). a. Effects of Strain Rate Reducing strain rate sharply reduces the ductility of Fe–24 at% Al tested in air at room temperature, as was discussed previously for Ni3Al (Fig. 2) and FeAl alloys (60). Such behavior further supports the concept of environmentally induced embrittlement, because lower strain rates provide more time for release of hydrogen from the water vapor in air, and for penetration into the alloy. The influence of strain rate on ductility of Fe–Al–Cr alloys in air has been confirmed by Shan and Lin (55). Further, Shea et al. (44) have shown the same trend of increasing embrittlement with lower strain rate in hydrogen-charged Fe–23.5 at% Al. Fig. 11 Effect of environment on room-temperature ductility of Fe–25 at% Al. (From Ref. 44.) 294 Stoloff b. Effects of Composition and Microstructure A fundamental question with respect to environmental embrittlement of Fe3Al alloys relates to the role of composition and, to lesser extent, the type and degree of long-range order. Early work on Fe–Al alloys showed that alloys with less than 8.5 wt% Al (16.3 at% Al) exhibited substantial ductility at room temperature (62,63). Recently, Vyos et al. (64) have shown that binary Fe–16.3 at% Al, which is a singlephase disordered alloy, is not susceptible to environmental embrittlement at 25°C. Microstructural changes seemed to have no effect on ductility, which was near 25% for the 16.3 at% Al alloy in the fully recrystallized condition. In contrast, an alloy containing 22 at% Al (displaying a two-phase ordered ⫹ disordered microstructure) showed tensile elongation of about 5% in air, with slightly higher elongations in vacuum. Unfortunately, the 16.3 at% Al alloy displays very high fatigue crack growth rates in air, suggesting that environmental embrittlement may be exacerbated under cyclic loading (3). With respect to chromium content, beneficial effects of chromium on environmental embrittlement in tension of Fe–28% Al alloys have been documented (61). The highest ductilities are produced in a partially recrystallized, quenched condition which produces partial B2-type order. c. Cyclic Loading Unlike the case for tensile ductility, under cyclic loading a marked susceptibility of a Fe–28 at% Al–4.8 Cr–0.47 Nb–0.2 C alloy (FA129) to environmental embrittlement has been noted, as shown in Figs. 12a and 12b (65). The effect is small in the B2 condition, but it is significant in the slow- Fig. 12 Effects of environment on fatigue crack growth of Fe3Al alloy FA-129 at room temperature. (From Ref. 65.) Embrittlement of Ni- and Fe-Based Intermetallics 295 cooled DO3 condition. Embrittlement does not seem to be influenced by the microstructure, as both fully recrystallized and partly recrystallized samples showed similar crack growth rates. The differing behavior between monotonic and cyclic loading may be due to the repeated rupturing of oxide films that occurs under cyclic loading. Therefore, any beneficial effect that chromium confers on the protectiveness of Al2O3 under monotonic conditions is lost under cyclic loading. Other solutes also have an effect. For example, the addition of 0.5% Zr to Fe–28 Al–5 Cr results in a decrease in the crack growth rate as compared with other alloys; see Fig. 13 (28). Carbon additions have been found to increase the critical stress intensity with little effect on the crack growth rate. There exists a limit to the beneficial effects of alloying with Zr, as a 1% Zr alloy has a higher crack growth rate than either of the 0.5% Zr alloys. The fatigue crack growth behavior of the iron aluminides demonstrates typical corrosion fatigue characteristics. A more inert environment results in a lower crack growth rate. In the alloys studied, the environment was found to have an effect on the threshold and critical stress intensities, except in the Zrcontaining alloys. In those alloys, the threshold stress intensity was found to be Fig. 13 Fatigue crack growth of iron aluminides in air at 25°C. Note low growth rates in alloys containing Zr. (From Ref. 28.) 296 Stoloff insensitive to environment. (These values are for comparison purposes only, as ASTM E-399 conditions for a valid fracture toughness test were not met; such values would probably be considerably lower than the ∆Kc values recorded.) d. Stress-Corrosion Cracking Alloy FA-129, which contains 28Al, 5 Cr, and 0.5% Nb, also has been shown to be susceptible to intergranular stress-corrosion cracking (SCC) in neutral-pH chloride-containing solutions (66). However, little effect of cathodic potential on SCC of Fe3Al was noted, in marked contrast to the high susceptibility of a Ti–Al–Nb–V–Mo alloy under similar charging conditions. These results conflict with the earlier findings of Shea et al. (44) and Kasul and Heldt (67), who showed that binary Fe–24.6 at% Al is very susceptible to hydrogen charging (see Fig. 14) as well as work that showed that the maximum stress intensity at fracture of cyclically loaded FA-129 is very sensitive to the atmosphere (65). These differences may be due in part to pH effects, as Ricker et al. (36) have reported that hydrogen-induced SCC of Ni3Al occurs only in lowpH solutions. In the case of Fe–24.6 at% Al ductility is reduced under cathodic conditions in both acidic and basic solutions. Buchanan et al. (68) also have reported on hydrogen-induced stress corrosion cracking of iron aluminides. D. Other Intermetallics There have been few reports of embrittlement of ordered iron- or nickel-based intermetallics by hydrogen or moisture, apart from those systems described in Secs. II.A and II.B. The influence of hydrogen charging on the tensile ductility of Ni2Cr was found to be a maximum for disordered and nearly ordered Ni2Cr (69). Recent fatigue-crack-growth studies have shown that nearly ordered Ni2Cr displays the lowest crack growth rate and highest threshold in air and in 3.5% NaCl, whereas disordered Ni2Cr had the highest growth rates and lowest threshold (70). Better fatigue-crack-growth resistance in partially ordered ane nearly ordered Ni2Cr was attributed to the occurrence of planar slip and the resultant cyclic-slip reversibility. However, cracks grow much more rapidly and thresholds are significantly lower in 3.5% NaCl than in air for all conditions of order (S ⫽ 0.5 or 0.9). Limited embrittlement studies have been carried out on the B2 alloy FeCo– 2% V (15). This alloy can be disordered by quenching from above the critical temperature, Tc, of 720°C. Hydrogen embrittlement occurs in the fully ordered (S ⫽ 1), partially disordered (S ⫽ 0.4), and disordered (S ⫽ 0) conditions. Brittle transgranular cleavage is observed in ordered material in both air and in hydrogen. Dimpled fracture is observed in partially ordered and disordered samples tested in air, whereas cleavage is again observed in hydrogen. These results support the conclusion that transgranular cleavage can be induced by hydrogen in B2 polycrystals (FeAl, Fe3Al, and FeCo–2% V) as well as in L12 single crystals Embrittlement of Ni- and Fe-Based Intermetallics 297 (Ni3Al,Ti; Co3Ti). The influence of hydrogen on cleavage energy of these cubic alloys needs to be established by both theoretical and experimental methods. Theoretical calculations have been provided thus far only for the influence of hydrogen on cleavage energy of FeAl (71), as described in the next subsection. E. Embrittlement Mechanisms As mentioned in the previous sections, many ordered intermetallics display environmental embrittlement in hydrogen (charged or uncharged) environments or in moist air at ambient temperatures. It has been shown that some intermetallics containing reactive elements, such as FeAl and Fe3Al, exhibit more severe embrittlement in moist air than in dry hydrogen. The proposed chemical reaction for the moisture-induced embrittlement of aluminides is (1) 2 M ⫹ 3 H2O → Al2O3 ⫹ 6 H (1) An alternative reaction between H2O and Ni3Al,Ti, as revealed by x-ray photoelectron microscopy (XPS) is (72) Al ⫹ x H2O → Al(OH)x ⫹ x H (2) It is the high-fugacity atomic hydrogen that rapidly penetrates into crack tips and causes severe embrittlement. For the B2 structure, the {200} planes offer the maximum amount of aluminum to react with water vapor. Fracture in FeAl is by transgranular cleavage in moisture or hydrogen, mixed mode in vacuum, and intergranular in dry oxygen. The underlying mechanism of moisture-induced embrittlement in FeAl, Fe3Al, Ni3Al, Co3V, and other intermetallics is undoubtedly embrittlement by hydrogen, with the principal difference being the manner in which atomic hydrogen is generated and absorbed at crack tips. The reaction kinetics of D2O molecules with single crystals of Ni3(Al,Ti) have been shown to strongly depend on the crystallographic orientation of the surface (73). Chemically adsorbed D2O reacts with {100} planes and generates deuterium when heated to at least 200 K. On the other hand, no deuterium was detected on {111} surfaces. Therefore, the decomposition of moisture depends on the atomic arrangement and chemical composition of crystallographic planes. Atomic hydrogen also has been detected from the reaction of moisture with an iron aluminide by laser desorption mass spectrometry (74). Cohron et al. (23) have shown that when low-pressure hydrogen is disociated in contact with an ionization gauge, the embrittlement of Ni3Al is severe. However, when the gauge is turned off, low-pressure hydrogen is not very embrittling. These results demonstrate that the dissociation of molecular hydrogen is a necessary precursor to severe embrittlement. The yield strength of the intermetallics is found to be insensitive to the test environment, as is typical with conventional alloys. Strain-rate effects are very significant, due to the time dependence of hydrogen diffusion. 298 Stoloff The highest ductility is generally obtained in dry-oxygen environments because oxygen suppresses the reaction of Eq. (1) and allows more rapid formation of oxides than in the presence of moisture. Hydrogen embrittlement is a very complex phenomenon even in conventional metals and alloys. The underlying mechanisms suggested for hydrogen embrittlement in ordered intermetallics resemble those for structural alloys and can be grouped into four categories: 1. 2. 3. 4. Reduction of atomic bonding across cleavage planes Reduction of cohesive strength across grain boundaries Reduction of dislocation mobility and crack tip plasticity Formation of brittle hydrides Environmental and hydrogen embrittlement in the bcc-ordered iron aluminides and in FeCo–V occurs along cleavage planes rather than grain boundaries, suggesting that cleavage strength is reduced by absorbed or adsorbed hydrogen. Experimental results are supported by the previously cited first-principles quantum-mechanical calculations, which indicate that absorbed hydrogen significantly reduces the cleavage strength and energy of FeAl (by as much as 20–70%, depending on the hydrogen concentration) (71). Superdislocations have been suggested to be the carriers for enhanced diffusion of hydrogen at crack tips (21). For face-centered-cubic (fcc)-ordered intermetallics, hydrogen embrittlement usually causes intergranular fracture, suggesting that reduction of cohesive strength along the boundaries is responsible. However, single crystals of various L12 alloys also are embrittled (8) so that grain boundaries clearly are not necessary to observe reduced ductility. Bond et al. (75) have shown by in situ transmission electron spectroscopy (TEM) that hydrogen enhances dislocation mobility and crack growth in Ni3Al, but these authors also suggested that decohesion is the failure mechanism. More work is needed to determine whether enhanced dislocation mobility in the presence of hydrogen occurs in other intermetallic systems. III. EMBRITTLEMENT AT ELEVATED TEMPERATURES Environmental degradation in iron- and nickel-based ordered intermetallics occurs also at elevated temperatures; see Table 6 (2). Hydrogen is the major embrittling agent and oxygen is beneficial at room temperature; oxygen is the major embrittling agent at elevated temperatures (typically above 300°C). At present, only a few intermetallic systems have been studied for environmental degradation at elevated temperatures and data are available principally for Ni3Al and Ni3Si alloys. Embrittlement of Ni- and Fe-Based Intermetallics Table 6 299 Elevated-Temperature Embrittlement of Ordered Alloys Alloy Crystal structure Environmental embrittlementa Ref. Ni3Al ⫹ Hf ⫹ B Ni3Al ⫹ Cr (Ni,Co)3Al Ni3Si Ni3Si ⫹ Cr Ni3(Si,Ti) ⫹ B (Fe22Co78)3V Ni3(Si,Ti) ⫹ B FeAl Fe3Al L12 L12 L12 L12 L12 L12 L12 L12 B2 DO3 䊊 䊊 䊊 䊊 䊊 䊊b 䊊 䊊 ⫻ ⫻ 76 76 93 80 80 38 81 38 95 95 䊊: Environmental embrittlement is observed when tested in oxidizing environments; ⫻: not observed. b No difference in elevated-temperature tensile ductility between air and vacuum, but environmental embrittlement is possibly masked because of a poor vacuum. Observed but reduced by alloying with Cr. Source: Ref. 2. a A. L12 Alloys 1. Ni3Al Tensile properties of Ni3Al are sensitive to test temperature and environment. Figure 14 compares the tensile elongation of a Ni3Al alloy (Ni–21.5 Al–0.5 Hf– 0.1 B) (IC-145) tested in air and vacuum (10⫺3 Pa) as a function of test temperature (76). The alloy tested in air showed appreciably lower ductility than that tested in vacuum at temperatures above 300°C, and the severest embrittlement occurred near 750°C, despite the fact that Ni3Al alloys exhibited good oxidation resistance in air. The loss in ductility generally is accompanied by a change in fracture mode from ductile transgranular to brittle intergranular. Similar embrittlement has been observed in other Ni3Al alloys, such as B-doped Ni3Al containing iron or hafnium (77). In these cases, oxygen has been identified as the embrittling agent. The similarity in the shape of the two curves in Fig. 14 further indicates that embrittlement cannot be completely suppressed by a conventional vacuum of 10⫺3 Pa. The role of air pressure in embrittlement of Ni3Al alloys has been shown clearly for IC-136 (Ni–23 at% Al–0.5% Hf–0.07% B) tested at 760°C. There is a rapid increase in elongation as air pressure is reduced from 1 to 10⫺3 torr, and continued to increase in ductility to a pressure of 10⫺7 torr (78). 300 Stoloff Fig. 14 Tensile ductility of IC-145 (Ni–21.3 at% A1–0.5 Hf–0.1 B) in air and in vacuum. (From Ref. 76.) Fortunately, alloying with 6–8 wt% Cr considerably diminishes the degree of embrittlement (76). Test environments also influence the fatigue life of boron-doped Ni3Al (24 at% Al) at elevated temperatures (79). The alloy showed a sharp drop in fatigue life at temperatures above 500°C when tested in a conventional vacuum (10⫺3 Pa); see Fig. 15 (80). The decrease in the fatigue life was accompanied by a change in fracture mode from transgranular to intergranular. 2. Ni3Si Alloys As in the case of Ni3Al, Ni3Si alloys exhibit severe environmental embrittlement in oxidizing environments at elevated temperatures. For Ni3Si and Ni3(Si,Ti) alloys doped with and without boron, tensile ductility decreases sharply at temperatures above 300°C in moist air. The ductility of low-Si alloys exhibits a minimum at 600°C, as seen in Fig. 16 (79); above that temperature, the ductility increases sharply. Ductility in vacuum at 600°C increases by a factor of 20. As in the case of Ni3Al, Cr additions in the range 2–6 at% effectively reduce embrittlemen (81). For high-(Si⫹Ti) alloys (e.g., 21%) ductility decreases continuously with increasing temperature and approaches zero at temperatures above 600°C. No improvement in ductility results in a conventional vacuum at these temperatures. Embrittlement of Ni- and Fe-Based Intermetallics 301 Fig. 15 Influence of temperature on fatigue life, Nf , of Ni3Al⫹B, tested in vacuum. (From Ref. 80.) B. Other Intermetallics (CoFe)3V alloys, which are susceptible to embrittlement by water or moist air (9), also display a sharp loss in ductility at temperatures above 500°C. The minimum ductility in both air and vacuum (⬃5 ⫻ 10⫺4 Pa) is noted near Tc (⬃910°C) (81). Ductility is lower in air than in vacuum between 500°C and 910°C. Embrittlement below Tc was attributed to oxygen-induced penetration of grain boundaries. These alloys are not embrittled at temperatures above Tc, probably due simply to the lack of long-range order. Iron aluminide alloys based on Fe3Al or FeAl do not appear to be susceptible to elevated-temperature embrittlement in oxidizing environments, in spite of their extreme ambient-temperature embrittlement in moist air. The absence of elevated-temperature embrittlement in moist air. The absence of elevated-temperature embrittlement in these aluminides is possibly related to the lack of a substantial yield anomaly, together with rapid formation of protective oxide films 302 Stoloff Fig. 16 Yield strength and elongation of Ni–18.9 at% Si, tested in air as a function of temperature. (From Ref. 79.) due to rapid diffusion. Further experimental studies are required to clarify these points. C. Embrittling Mechanisms Embrittlement (other than via the pest reaction) has been suggested to be caused by a dynamic effect simultaneously involving high localized-stress concentration, elevated temperature, and gaseous oxygen (2). Such a dynamic effect involves repeated weakening and cracking of grain boundaries as a result of oxygen absorption and penetration at crack tips. Based on a detailed study of crack growth in Ni3Al alloys tested in oxidizing environments, a fracture mechanism of stressassisted grain-boundary oxygen penetration has been suggested to explain the elevated-temperature embrittlement (82). This model consists of four sequential steps: (1) occurrence of gaseous oxygen to the crack tips where a high localized stress field is involved, (3) oxygen penetration on its atomistic form to the stress field ahead of tips, and (4) inward development of surface cracks preferentially along the grain boundaries, leaving some secondary cracks (83). Steps (2) and (4) proceed continuously and repeatedly during deformation, leading to premature Embrittlement of Ni- and Fe-Based Intermetallics 303 fracture and severe loss in ductility at elevated temperatures in oxidizing environments. The fatigue crack growth rate of Ni3Al and the Ni3Al-based alloy IC-221 increases with increasing temperature, between 25°C and 600°C, even when testing is carried out in moderate vacuum (83). This behavior indicates that environmental embrittlement due to oxygen is occurring, especially because the flow stress increases with temperature over the same range. These observations were consistent with those of Hippsley and Devan (82) for static crack growth. By contrast, when Fe3Al is tested at elevated temperatures in air the fatigue crack growth rate decreases, perhaps because moisture-induced embrittlement is maximized near room temperature, as in the case of structural steels. Alternative mechanisms for high-temperature embrittlement also have been suggested by observations of low elevated-temperature ductility in Ni3Al-based single crystals tested in air (84–86). The increase in flow stress to a maximum in the temperature regime of minimum ductility may play a significant role in elevated-temperature embrittlement. Also, Yizhang et al. (87) suggest that in cast polycrystalline Ni3Al alloys, chromium may enhance ductility through its effect on γ-γ′ eutectic size and distribution and the replacement of brittle γ′–γ′ grain boundaries by tough γ′–γ–γ ′ boundaries. Clearly, the role of chromium in improving ductility in air requires further study. IV. ALLEVIATION OF EMBRITTLEMENT Environmental degradation has been identified as a main cause of the low ductility and brittle fracture in many ordered intermetallics. This problem has to be solved satisfactorily in order to use intermetallic alloys as engineering materials. Despite their different embrittling agents, ambient-temperature and elevated-temperature embrittlement can be treated together because both involve surface reactions and are sensitive to localized-stress concentrations. Results generated to date indicate that embrittlement can be alleviated or reduced by (a) control of surface conditions, (b) control of grain size and shape, (c) alloy additions, (d) processing techniques, and (e) prestrain. A. Surface Conditions Control of surface conditions is a straightforward way to alleviate environmental degradation involving surface reactions. In several cases, preoxidation and formation of protective oxide scales were proven to be beneficial in reducing environmental embrittlement at ambient and elevated temperatures. Preoxidation at 1000°C effectively reduced ambient-temperature embrittlement in B-doped Ni3Al charged with hydrogen (15). Formation of protective oxide scales increases the 304 Stoloff tensile ductility of FeAl and Fe3Al alloys in air at ambient temperatures and the ductility of B-doped Ni3Al alloys at elevated temperatures (76). Unfortunately, the oxide films crack after stretching a few percent and their protective effect disappears. Surface coatings also should be useful in protecting underlying alloys from hydrogen or oxygen penetration along grain boundaries or through the lattice; however, this effect has not yet been well demonstrated. B. Grain Size and Shape Columnar-grained structures have proven to be effective in reducing environmental embrittlement in Ni3Al alloys of varying aluminum contents, tested in moist air at room and elevated temperatures (88–90). For example, formation of a columnar-grained structure in boron-doped Ni3Al produced by directional levitation-zone remelting increases the ductility in air from 0.2% to 33% at temperatures in the range 600–700°C (88). The loss of ductility in air is accompanied by a change in fracture mode from microvoid coalescence to intergranular in equiaxed material. However, the columnar structure displayed mainly transgranular failure in both air and vacuum. The beneficial effect of the columnar-grained structure with grain boundaries oriented parallel to the stress axis is attributed to minimizing the normal stress across grain boundaries, thereby suppressing nucleation and propagation of cracks along boundaries even when those boundaries are weakened by oxygen penetration. Hirano (89) has reported high ductility in stoichiometric Ni3Al tested both parallel and perpendicular to the growth direction. More recent work shows that elongation increases with decreasing solidification rate; see Fig. 17 (90). The technique also is effective in Al-rich Ni3Al, leading the authors to suggest that unidirectional solidification may be effective in improving ductility of two-phase Ni3Al–NiAl alloys. A detailed analysis of grain-boundary chemistry and misorientation is required to fully understand these results. However, it appears that the mostly low-angle and ⌺3-type boundaries resulting from unidirectional solidification are less affected by moisture than are general grain boundaries in conventionally processed material (91). Studies of unidirectional solidification should be extended to other intermetallic systems. The role of grain size seems to be the same for intergranular and transgranular fracture in that refining grain size tends to reduce susceptibility to embrittlement both at room and elevated temperatures. Takeyama and Liu (92) have shown that heat treatment of Ni3Al alloys in oxidizing atmospheres can cause embrittlement in subsequent tests at both ambient and elevated temperatures. The degree of embrittlement is essentially zero for fine-grained (20 µm) alloys, but it is very pronounced at 200 µm, as shown in Fig. 18 (92). Apparently, a thin protective Al-rich oxide film is formed on fine-grained material, whereas a less protective, predominantly Ni-rich film forms on coarse-grained samples. The aluminum-rich oxide on fine-grained material may form as a result of short-circuit diffusion of aluminum atoms from the interior to the surface. Embrittlement of Ni- and Fe-Based Intermetallics 305 Fig. 17 Effect of growth rates (mm/h) on ductility of Ni–26 at% Al. (From Ref. 90.) Fig. 18 Effect of grain size on ductility of bare and preoxidized Ni–23 at% Al–0.5 Hf– 0.5 B at 600°C and 760°C in vacuum. (From Ref. 92.) 306 Stoloff C. Alloy Additions Several striking examples of alleviation of embrittlement by alloying may be cited. The ductilizing effect of boron on low-temperature fracture of Ni3Al and Ni3Si in air is clearly associated with the inhibition of moisture-induced embrittlement (22,23). Similarly, chromium enhances high-temperature ductility of Ni3Al, presumably by altering the kinetics of oxide film formation (76). Chromium also has a beneficial effect on the low-temperature ductility of Fe3Al alloys in air (61). However, chromium does not prevent environmental embrittlement under cyclic loading conditions (44,60). Finally, zirconium additions have been shown to improve both ductility and fatigue crack growth behavior in air of Fe–28 Al–5 Cr alloys, although the mechanism for such improvements remains obscure (28). D. Processing Techniques Boron-doped Ni3Al alloys prepared by rapid solidification are more prone to environmental embrittlement than alloys produced by conventional casting and thermomechanical processing (93). The latter exhibit much higher ductility at 750°C in vacuum than does the rapidly solidified material, although at 600°C, there is little effect of microstructure. The same workers reported that spray-formed Ni– 10 at% Co–24% Al with boron exhibited zero ductility in vacuum at 760°C, but ductility could be substantially improved by subsequent thermomechanical processing. A number of factors could be influencing the results obtained with differing processing techniques, including grain size and shape, grain-boundary energy, oxygen and other impurities, mobile dislocation density, thermal history, and texture. Of these, the only factor definitely linked with low ductility in many alloys is impurity content. E. Prestrain There is now considerable evidence that prestrain can reduce the severity of environmental embrittlement in Ni3Al and Ni3Si alloys (35,40,94). The magnitude of the effect depends on the degree of prestrain as well as the temperature at which it is performed. It remains to be seen whether this effect also can be found in other intermetallics, especially the iron aluminides. There are several possible mechanisms for the prestrain effect, such as dislocations acting as traps for hydrogen or the creation of compressive residual stresses in the case of prestrain by shot peening. It has been argued that the beneficial effects of prestrain must mean that dislocation-assisted transport of hydrogen does not occur. However, it must be pointed out that the dislocation transport mechanism depends on the simultaneous application of strain and exposure to hydrogen, a condition clearly unmet by the prestrain experiments. Embrittlement of Ni- and Fe-Based Intermetallics V. 307 SUMMARY Many nickel- and iron-based intermetallics are embrittled by moisture and other hydrogen-containing environments at low temperatures and by oxygen at elevated temperatures. Crystal structure, composition, metallurgical variables, and test conditions all play a role in determining the degree of embrittlement. 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INTRODUCTION Nonoxide ceramics such as silicon carbide (SiC) and silicon nitride (Si3N4) are promising materials for a wide range of high-temperature applications. These include such diverse applications as components for heat engines, high-temperature electronics, and reentry shields for space vehicles. Table 1 lists a number of selected applications. Most of the emphasis here will be on SiC and Si3N4. Where appropriate, other nonoxide materials such as aluminum nitride (AlN) and boron nitride (BN) will be discussed. Proposed materials include both monolithic ceramics and composites. Composites are treated in more detail elsewhere in this volume; however, many of the oxidation/corrosion reactions discussed here can be extended to composites. In application, these materials will be exposed to a wide variety of environments. Table 1 also lists reactive components of these environments. It is well known that SiC and Si3N4 retain their strength to high temperatures. Thus, these materials have been proposed for a variety of hot-gas-path components in combustion applications. These include heat-exchanger tubes, combustor liners, and porous filters for coal combustion products. All combustion gases contain CO2, CO, H2, H2O, O2, and N2. The exact gas composition is dependent on the fuel-to-air ratio or equivalence ratio. The equivalence ratio (EQ) is 311 Selected Applications Which Are Proposed for Nonoxide Ceramics Nonoxide ceramic Application Turbine engine components Combuster liners Blades and vanes Piston engine components Pistons Valves Industrial furnaces Heat exchangers Coal combustion Particulate filters Chemical process vessels, coal gasifiers, waste incinerators Reentry shields 312 Table 1 Approximate use temperatures (°C) SiC, Si3N4 Composites 900–1400 SiC, Si3N4 Composites 900–1400 SiC Composites SiC Composites SiC, Si3N4 900–1400 700–1000 900–1400 Electronic substrates AlN Fiber Coatings for composites BN Crucibles, insulators Liquid metal containers Processing, heat transfer Pump bearings, cooling lines for nuclear factors BN Various carbides and nitrides 900–1400 600–1400 SiC, Si3N4 300–600 1000–1500 Use: 600 Processing: 1200 Use: 200 Processing: 1000 900–1400 Ref. Combustion gases; deposits: Na, Mg, Ca sulfate, sodium vanadates Combustion gases 1 Combustion gases, various deposits Combustion gases, slag deposits 3 Various gases including air, H2S, HCl Reduced-pressure N2,O2, CO2, N, O Air Air Reduced-pressure combustion gases Vacuum, inert gases Vacuum or inert gas and liquid metals High-pressure fluids, 10–100 atm 2 4 5 6, 7 8 9 10 11 12 13 Jacobson and Opila High-temperature semiconductors SiC Composites SiC Environment Nonoxide Ceramics 313 a fuel-to-air ratio, with total hydrocarbon content normalized to the amount of O2 and defined by EQ ⫽ 1 for complete combustion to CO2 and H2O.] Figure 1 is a plot of equilibrium gas composition versus equivalence ratio. Note that as a general rule, all combustion atmospheres are about 10% water vapor and 10% CO2. The amounts of CO, H2, and O2 are highly dependent on equivalence ratio. Other proposed applications for SiC include high-temperature semiconductors and for AlN include electronic substrates. In these situations, high-temperature oxidation behavior is a prime issue. Reentry shields have long been a useful Fig. 1 Equilibrium gas composition versus equivalence ratio. (From Ref. 24.) 314 Jacobson and Opila application of ceramic materials—here, the environment is a complex mixture of atoms and molecules. Finally, a growing area is the application of ceramics as pump components and ball bearings. The environment depends on the fluid of interest, but it is generally high-temperature high-pressure water. In this chapter we will discuss the interaction of nonoxide ceramics with some representative environments. We begin with pure oxidation, which is important for nearly all the applications, and then proceed to the more systemspecific environments such as molten salts and high-pressure water. II. HIGH-TEMPERATURE OXIDATION Figure 2 summarizes the oxidation behavior for a range of nonoxide ceramics (14). These data are presented in terms of recession in 100 h. Clearly, longterm operation requires very low recession. The key oxidation reactions can be summarized as follows: 1. Silica formers: SiC, Si3N4, MoSi2 MSi ⫹ 3/2O2(g) ⫽ SiO2 ⫹ MO (1) The oxidation of SiC and Si3N4 is described in detail in Sec. II.A. 2. Alumina formers: primarily alloys, but also AlN 2Al ⫹ 3/2O2(g) ⫽ Al2O3 (2) The underline indicates that aluminum is at less than unity activity, as would be found in an alloy. 2AlN ⫹ 3/2O2(g) ⫽ Al2O3 ⫹ N2(g) (3) The oxidation of AlN is discussed more fully in Sec. II.B. 3. Borides 2MB ⫹ 5/2O2(g) ⫽ B2O3(l) ⫹ 2MO (4) 4. Carbides MC ⫹ O2(g) ⫽ MO ⫹ CO(g) (5) 5. Nitrides MN ⫹ O2(g) ⫽ MO ⫹ NO(g) (6) or 2MN ⫹ O2(g) ⫽ 2MO ⫹ N2(g) (7) Nonoxide Ceramics 315 Fig. 2 Recession due to oxidation for selected materials, with protective oxide scale indicated. (From Ref. 14.) Most of these materials oxidize according to a parabolic rate law, but some oxidize according to linear kinetics. Hence, Fig. 2 gives oxidation kinetics in terms of recession to account for both kinetic laws. The reader is referred to Ref. 14 and references contained therein for more details on the kinetics of these reactions. The important point from Fig. 2 is that only materials which form protective silica or alumina films are useful for long times in oxidizing environments. Most of the borides, carbides, and nitrides do not form protective metal oxide scales. Hence, the focus of this chapter will be on silica-forming ceramics, SiC and Si3N4, and the alumina-forming ceramic, AlN. We shall also include a brief discussion of BN, which is an important material in many areas of technology. A. Oxidation of SiC and Si3N4 1. Oxygen Transport in Silica The low oxidation rates of silica-forming ceramics are due to remarkably low oxygen transport rates in silica. Thus, we begin with a brief discussion of oxygen 316 Jacobson and Opila transport in silica. Amorphous silica, found for short-time oxidation or in very clean systems, consists of a random network of Si–O tetrahedra as shown schematically in Fig. 3a (15). Cristobalite, the crystalline form of silica most often observed after long-term oxidation or oxidation in less clean environments, is composed of the same Si–O tetrahedra, but arranged in an orderly structure as shown in Fig. 3b. Oxygen transport through these structures occurs by two different mechanisms. First, molecular oxygen can move through the interstices between the tetrahedra by a permeation mechanism. Alternatively, ionic oxygen can move from network site to network site by a bond-breaking exchange mechanism. The permeation mechanism requires less activation energy because bond breaking is not required. Permeation through cristobalite is expected to be slower than through amorphous silica because the regular structure is restricted to six-member oxygen rings, whereas the irregular structure found in amorphous silica allows for seven- and eight-member rings (15). Permeation rates of molecular oxygen in amorphous silica have been determined (16); however, the corresponding rates in cristobalite have not been measured. The measurements are difficult because bulk cristobalite is not available and must be nucleated and characterized at temperature. In addition, the presence of short-circuit transport paths along cristobalite grain boundaries and cracks formed from the β to α cristobalite phase transformation upon cooling make the measurement of intrinsic oxygen transport rates in cristobalite impossible. 2. Oxidation of Silica-Forming Ceramics: Experimental Techniques Oxidation kinetics for these materials are determined by measuring either weight changes using a sensitive thermogravimetric balance as shown in Fig. 4 or oxide thickness changes using optical techniques. An example of weight change versus time for SiC oxidation in dry oxygen is shown in Fig. 5. Because the oxidation rate is controlled by diffusion of oxygen through the silica scale, the weight change decreases with time according to a parabolic law and is described by the parabolic rate constant kp. Accurate determinations of oxidation kinetics are more difficult for silica-forming materials than for other metal-oxide-forming materials because silica grows at such a slow rate and because silica growth kinetics are affected by even small amounts of impurities in the sample or oxidation environment (17). A comparison of the parabolic oxidation rates for silicon, SiC, and Si3N4 as a function of temperature is shown in Fig. 6. Similarities and differences in the oxidation kinetics of these materials are discussed in Secs. II.A.3–II.A.6. It is instructive to begin with the relatively simple case of silicon oxidation. We will then focus on very pure SiC and Si3N4 in order to understand the Nonoxide Ceramics 317 Fig. 3 Structure of SiO2: (a) amorphous; (b) crystalline; (c) with sodium cations. (Adapted from Ref. 15.) 318 Fig. 4 Jacobson and Opila Schematic of electrobalance and vertical tube furnace used for oxidation studies. fundamental oxidation behavior of these materials. These pure materials are also important in microelectronic applications. 3. Oxidation of Silicon Silicon oxidation has been studied in great detail due to its application in microelectronic devices. The model of Deal and Grove (18) has been used to successfully describe silicon oxidation behavior at all but the shortest times. This model considers two possible rate-limiting steps for the oxidation of silicon. First, the reaction of oxygen with silicon according to a linear rate law controls the oxidation rate at short times or for thin scales. Second, transport of oxygen through Nonoxide Ceramics 319 Fig. 5 Oxidation of SiC at 1300°C in 1 atm oxygen, showing parabolic kinetics. (Adapted from Ref. 20.) 320 Jacobson and Opila Fig. 6 Arrhenius plot of log kp versus 1/T obtained in dry oxygen for Si, chemicallyvapor-deposited (CVD) SiC, CVD Si3N4, and Si3N4 with additives. (Data from Refs. 18, 20, and 35.) the growing silica scale controls the oxidation rate according to a parabolic rate law for long times and for thick scales. These two rate laws are combined into a single expression that is valid for all times: 冢 冣 t⫹τ xo ⫽ 1⫹ 2 1/2 A A /4B 1/2 ⫺1 (8) where xo is the oxide thickness, t is time, τ is the offset time which corrects for the presence of an initial oxide layer, B is the parabolic rate constant, and B/A is the linear rate constant. At short times, this expression simplifies to the linear rate law xo ⫽ (t ⫹ τ)B/A; at long times, the expression simplifies to the parabolic law x 2o ⫽ Bt. For long-term applications of silica-forming materials, the parabolic law usually describes the oxidation kinetics adequately. Deal and Grove have derived the parabolic rate constant in terms of the properties of silica; that is, B⫽ 2DeffC* N (9) Nonoxide Ceramics 321 and C* ⫽ kP n (10) where Deff is the effective diffusivity of the oxidant in silica, C * is the equilibrium concentration of the oxidant in the oxide, N is the number of oxidant molecules incorporated into a unit volume of the oxide layer, P is the oxidant pressure, n is a power-law exponent, and k is a constant. The parabolic rate constant was found to vary linearly with the oxidant pressure for both oxygen and water vapor (i.e., n ⫽ 1), indicating that the oxidant did not dissociate and molecular permeation through the silica is the rate-limiting step for oxidation in the parabolic regime. 4. Oxidation of SiC The oxidation of SiC is similar to the oxidation of silicon because a silica scale forms in both cases. In this case, however, the gases generated, as shown in the following oxidation reactions, will cause some differences: SiC ⫹ 3/2O2(g) ⫽ SiO2 ⫹ CO(g) (11) SiC ⫹ 2O2(g) ⫽ SiO2 ⫹ CO2(g) (12) The linear reaction rate is different because the oxidation of C to CO(g) or CO2(g) occurs in addition to the oxidation of Si. The parabolic oxidation rate could be different if the outward transport of CO or CO2, rather than the inward transport of oxygen, limited the oxidation rate. It has been shown that the oxidation rates of SiC are about a factor of 2 slower than silicon (19–21) due to the extra consumption of oxygen in the reaction with C, as predicted (22). In addition, the activation energy for oxidation (20) is nearly identical to that of silicon (18) (see Fig. 6) and the permeation of molecular oxygen through silica (16). Finally, the oxidation rate of SiC is found to depend on the oxygen partial pressure (23). Therefore, it is generally agreed that oxygen transport inward is the rate-limiting step for parabolic oxidation of SiC. This issue has been discussed more fully in Refs. 24 and 25. The dependence of the parabolic rate constant on the oxygen partial pressure, given by the power-law exponent, n, in Eq. (10), gives information about the type of oxygen transport occurring in the silica. For n ⫽ 1, molecular permeation of oxygen occurs. For n ⬍ 1, some dissociation of oxygen into a charged species occurs and network diffusion of oxygen by a bond-breaking process is likely. Zheng et al. (23) have determined the power-law exponent for the oxidation of SiC to vary between 0.6 and 0.3 at temperatures from 1200°C to 1500°C. This implies that some combination of permeation and network diffusion limits the oxygen transport through the silica scale grown on SiC, with network diffusion increasing with temperature. This is confirmed by 18O tracer diffusion stud- 322 Jacobson and Opila ies, which show isotope exchange with network oxygen becomes increasingly important as the temperature is increased (26). However, the similar activation energies for molecular oxygen permeation of silica (16) and for oxidation of SiC up to temperatures of 1500°C (20) indicates that oxygen transport is dominated by the permeation mechanism. The effect of crystallization of silica on SiC oxidation is a complex topic. Some investigators have used crystallization to explain observed rate changes and/or activation energy changes. However, current evidence indicates that many of these deviations in rate law and activation energies can be explained by impurities in either the environment or the SiC (17,20,27,28). Ogbuji (29) has shown that silica scales fully crystallized during argon anneals at 1300°C do result in slower oxidation rates by about a factor of 30, but that fully crystalline scales are not found in actual practice because amorphous silica is continually formed during the course of SiC oxidation. 5. Oxidation of Si3N4 Like SiC, the oxidation of Si3N4 results in the formation of a silica layer and the generation of gaseous products according to the following simplified reaction: Si3N4 ⫹ 3O2(g) ⫽ 3SiO2 ⫹ 2N2(g) (13) An additional complication, however, is the formation of a suboxide layer of amorphous silicon oxynitride of variable stoichiometry (30,31). The oxidation reaction can be written as (32) Si3N4(1⫺x) O6x ⫹ 3δxO2(g) ⫽ Si3N4(1⫺x⫺dx) O6(x⫹dx) ⫹ 2δxN2(g) (14) for x varying between 0 and 1. Measurements of the oxidation kinetics of Si3N4 have shown they are parabolic but significantly slower than oxidation of both silicon (33) and SiC (20) to temperatures of 1500°C. Pure Si3N4 is the slowest oxidizing material known today. The higher activation energy for Si3N4 reflects the additional energy required in the breaking of bonds for the nitrogen–oxygen substitution reaction. The oxidation rates were found to be dependent on the oxygen partial pressure, but independent of the nitrogen partial pressure (33). In this case, the oxidation reaction is limited by oxygen transport and reaction in the oxynitride layer (32) rather than by oxygen transport in silica. 6. Oxidation of Additive-Containing Materials The above discussions have considered only very pure materials. SiC and Si3N4 materials used for structural applications often contain additives to aid sintering. These additives affect long-term oxidation behavior. First, additives can diffuse into the silica scale during oxidation (34–38). Impurities present in silica increase oxygen transport rates by modification of the silica network structure and thereby Nonoxide Ceramics 323 increase oxidation rates of the silica-forming material (39). Second, additives such as Y2O3, La2O3, MgO, and CeO2 present in significant amounts (4–10%) will diffuse to the silica surface and form discrete particles of silicates often of the M2Si2O7 phase where M ⫽ Y, La, Ce. In these cases, the oxidation kinetics are still parabolic to times as long as 1000 h (D.S. Fox, personal communication, 1997), but are limited by the outward diffusion of metal cations to the matrix– oxide interface (34,35). In addition, for silicon nitride ceramics, the presence of impurities (40) and additives (41) prevent the formation of the silicon oxynitride inner layer; thus, oxidation rates of Si3N4 in practical applications are never as low as those found for pure chemically-vapor-deposited (CVD) Si3N4. As additive levels and impurities increase further, the oxidation kinetics often do not follow simple parabolic rate laws. These more complex oxidation kinetics have been modeled in various ways (42,43). Finally, for systems containing large amounts of additives, ⬎10%, at intermediate temperatures of 1000–1200°C, reactions to form silicates occur in the grain boundaries with a volume expansion large enough to cause disintegration of the ceramic (44). 7. Thermal Cycling Effects Long-term applications of structural ceramics, such as heat engines, require thermal cycling. Thermal cycling is a concern for several reasons. First, upon cooling, cristobalite undergoes the β to α phase transformation with an accompanying 3% volume contraction (45). This volume change causes cracks to form in cristobalite upon cooling. Second, the thermal mismatch between SiC or Si3N4 and cristobalite, shown in Fig. 7, results in tensile stresses in the oxide layer upon cooling. Cracks in the oxide may then form. However, cyclic oxidation tests at 1300°C in 5-h cycles for 1000 h have shown few deleterious effects on the oxidation kinetics when measured by weight change (46), as shown in Fig. 8. One possible explanation is that upon reheating, stress is relieved and the cracks heal. Note that the tensile stresses formed in these ceramic–oxide systems contrast with those in superalloy–oxide systems where compressive stresses form in the scale on cooling and oxide spallation is typically observed. B. Oxidation of Other Nonoxide Ceramics BN is a useful crucible material in a vacuum or inert atmosphere. However, in oxygen, it readily forms a liquid B2O3 scale which is not protective (47). Furthermore, this B2O3 scale readily reacts with water vapor and forms stable volatile H–B–O(g) species, such as HBO2(g), H3BO3(g), and H3B3O6(g). These species have strongly negative free energies of formation and form even with parts-permillion (ppm) levels of water vapor. A similar situation exists with TiB2, which is an attractive structural material due to its high strength (48). 324 Jacobson and Opila Fig. 7 Thermal expansion of SiC, Si3N4, amorphous SiO2, and crystalline SiO2 as a function of temperature. (From Ref. 24.) AlN is currently a candidate for electronic substrates due to its high thermal conductivity (9). As an alumina former, it is expected to exhibit slow oxidation kinetics. However, it appears to oxidize quite rapidly above about 1000°C, with rates quite dependent on grain size, porosity, the presence of second phases, and impurity content (49). In many instances, linear kinetics have been reported, in contrast to alumina formation on metal alloys (50). In addition, it is well established that water vapor in the oxidizing stream leads to extensive attack (51). The reasons for this extensive alumina formation in dry and wet oxygen are not Nonoxide Ceramics 325 Fig. 8 Cyclic oxidation kinetics of SiC and Si3N4 obtained in air at 1300°C. (Adapted from Ref. 46.) clear. Clearly, nitrogen must escape through the alumina scale, which may lead to micropore formation. It has been suggested that the injection of N3⫺ leads to excess vacancy formation in the alumina and more rapid diffusion rates (52). In addition, the coefficient of thermal expansion mismatch between AlN and Al2O3 may lead to scale cracking (53). III. COMPLEX ENVIRONMENTS In most applications, ceramic materials are subjected to more aggressive environments than high-temperature oxygen alone. These are outlined in Table 1. As in the case of pure oxidation, most of the available data on interactions of nonoxide ceramics in complex gas mixtures are for SiC and Si3N4. A. Water Vapor There is general agreement that water vapor enhances the oxidation rate of silicon, SiC, and Si3N4. There is a large amount of disagreement as to the magnitude of this effect for SiC and Si3N4. This disagreement arises in part from the complex effects water vapor has on the growing silica scale. These effects include enhanced impurity transport to the silica scale in water vapor containing atmo- 326 Jacobson and Opila spheres, enhanced solubility of water in the silica scale, alterations in the silica scale viscosity and structure and, finally, formation of volatile silicon hydroxide and oxyhydroxide species. Each one of these topics will be discussed separately below. Again, references to the oxidation behavior of silicon in water vapor are made for comparison. Oxidation of silicon by water vapor occurs by the following reaction: Si ⫹ 2H2O(g) ⫽ SiO2 ⫹ 2H2(g) (15) The enhanced growth rates of SiO2 by this reaction, relative to dry oxygen, has commercial application in the more rapid fabrication of microelectronic devices. Oxidation of SiC by water vapor occurs by the reactions (54) SiC ⫹ 3H2O(g) ⫽ SiO2 ⫹ CO(g) ⫹ 3H2(g) for T ⬎ 1400K (16) SiC ⫹ 2H2O(g) ⫽ SiO2 ⫹ CH4(g) for T ⬍ 1400K (17) and for Si3N4 by the reaction Si3N4 ⫹ 6H2O(g) ⫽ 3SiO2 ⫹ 2N2(g) ⫹ 6H2(g) (18) Note that in each case, the solid product SiO2 is formed with the generation of additional gaseous products. 1. Impurity Transport Impurities normally found in water-vapor-containing environment form M– OH(g) species, where M is the impurity element such as Na, K, Fe, and so forth. Because M–OH(g) species are so thermodynamically stable, the quantity of impurities transported to a silica scale forming on SiC or Si3N4 in a water-vapor containing environment is increased. This increased contamination of the silica scale results in faster transport rates of oxidant through the scale, and thus increased oxidation rates of SiC or Si3N4. This effect has been identified for both Na (55) and K contamination (56) during the oxidation of SiC. 2. Enhanced Solubility of Water in the Silica Scale Deal and Grove (18) have shown that the parabolic oxidation rate of silicon in water vapor is increased by about an order of magnitude over the rate found in dry oxygen. This is explained by examination of Eq. (9). Although the diffusivity of water vapor in silica is almost two orders of magnitude slower than molecular oxygen, the solubility of water is nearly three orders of magnitude larger than oxygen. The net result is that parabolic oxidation rates of silicon in water vapor are more than one order of magnitude larger than those observed in dry oxygen. This explanation is directly applicable to the discussion for SiC and Si3N4 because silica is the oxidation product for these materials. This enhancement in parabolic Nonoxide Ceramics 327 oxidation rate constants has been observed for SiC (55,57) and in some cases for Si3N4 (58). Deal and Grove (18) have also demonstrated that the parabolic oxidation rate has a power-law dependence on the water-vapor partial pressure, with the power-law exponent, n [Eq. (10)], equal to 1, indicating that molecular water diffusion is the rate-controlling step in the oxidation of silicon. Power-law exponents for the oxidation of SiC of 0.67 (59) and 0.67 to 0.85 (54) have been obtained. This indicates dissociation of water into a singly-charged species is probable. The discrepancy between the findings for silicon and SiC is unexpected because the transport properties of silica should be independent of the substrate material. A complex relationship between parabolic rate constant and water-vapor partial pressure has been observed for Si3N4 (58), but this may be explained by silica volatility, as described in Sec. III.A.4. 3. Alterations in the Structure of Silica It has been shown that the viscosity of amorphous silica decreases as the hydroxyl content increases (60). Hydroxyl groups are effective in breaking SiEOESi bonds in amorphous silica. It is suggested that the resulting silica allows the more rapid permeation of molecular oxygen (57). The reduced viscosity of amorphous silica in conjunction with increased amounts of gaseous products results in the formation of bubbles in the scale formed on SiC (54), as shown in Fig. 9. These Fig. 9 Bubbles formed in SiO2 due to oxidation of SiC in 90% water vapor/10% O2 at 1200°C. (From Ref. 54.) 328 Jacobson and Opila bubbles, in turn, create shorter transport paths for the oxidant to the SiC/SiO2 interface by decreasing the effective oxide thickness, thereby increasing the oxidation rate. In contrast, bubbles are not observed in the scales formed on SiC in dry oxygen or on silicon in wet oxygen. 4. Silica Scale Volatility Water vapor reacts with the silica scale formed on SiC and Si3N4 to form volatile hydroxide and oxyhydroxide species by the following reactions: SiO2 ⫹ 2H2O(g) ⫽ Si(OH)4(g) SiO2 ⫹ H2O(g) ⫽ SiO(OH)2(g) (19) (20) These volatile silicon hydroxides and oxyhydroxides have been identified experimentally using the transpiration technique (61) as well as mass spectrometry (62). The reaction of water vapor with SiC or Si3N4 involves the oxidation reaction [Eqs. (16)–(18)] and the simultaneous linear volatilization reaction [Eqs. (19) and (20)] resulting in overall paralinear kinetics (63). Paralinear kinetics, as measured by weight change, are shown in Fig. 10. At long times, the weight change and recession can be approximated by the linear volatilization rate alone. At this time, a steady state is achieved—the silica scale is consumed at the same rate it is formed, leaving a constant oxide thickness. The volatilization rate of silica is controlled by transport of the volatile species through a gaseous boundary layer Fig. 10 Schematic showing components of paralinear kinetics. (From Ref. 63.) Nonoxide Ceramics 329 (63). This boundary-layer-controlled volatility rate can be expressed in terms of the application conditions such as gas velocity, total pressure, and water-vapor partial pressure. For Si(OH)4(g) formation, the volatility rate has the following dependences: kl ⬀ v 1/2P(H2O)2 1/2 P total (21) where kl is the linear volatility rate, v is the gas velocity, and P is pressure. Thus, in high-pressure and high-velocity applications, such as a gas turbine engine, the recession of a SiC or Si3N4 component by the volatility mechanism can be significant. In summary, water vapor has deleterious effects on the durability of silica formers through a number of different mechanisms: increased transport of impurities to the oxide surface, increased oxide formation rates due to impurity effects, enhanced solubility of water vapor, short-circuit paths for oxidant transport, increased permeability of the oxidant, and, finally, consumption of the silica scale and component recession due to silica volatility. B. Carbon Dioxide Although CO2 is a major component of combustion gases (Fig. 1), there are only a few limited oxidation studies of SiC in CO2 (64–67). It has been found that the oxidation rates of SiC in CO2 are less than those in oxygen. Because the oxide growth rates are so low, it is difficult to determine whether the kinetics are linear, parabolic, or more complex. The oxidation weight gain for SiC in CO2 is shown compared to that observed in oxygen and a 50% water vapor–oxygen mixture in Fig. 11 (67). Thus, in a complex combustion environment, the effects of CO2 as an oxidant are negligible. C. Effects of Low P(O2), Reducing Gases, H2S, Cl2 1. SiO(g) Formation A unique, but very important issue with silica formers is the highly stable volatile suboxide, SiO(g). Consider the free energy of formation at 1500 K: /2Si ⫹ 1/2O2(g) ⫽ 1/2SiO2(s), ∆G ⫽ ⫺322 kJ/mol (22) ∆G ⫽ ⫺227 kJ/mol (23) 1 Si ⫹ /2O2(g) ⫽ SiO(g), 1 Note that these equations are normalized to 1 mol of oxygen atoms. The free energies of formations are close, indicating that SiO(g) can readily form. There are two conditions which lead to SiO(g) formation (68): active oxidation and oxidation in mixed oxidizing/reducing gases. In the active oxidation 330 Jacobson and Opila Fig. 11 Comparison of SiC oxidation in 50% H2O/O2, O2, and CO2 at 1200°C. (From Ref. 67.) case, the partial pressure of oxygen is too low to form a stable SiO2 film, but sufficient to form SiO(g). This can occur in certain heat-treating environments (69). Let us begin with a bare SiC surface. As the partial pressure of oxygen is increased, SiO(g) will form in increasing quantities. Then, SiO(g) formation will stop and a stable SiO2 film will form. This is the active-to-passive transition. Wagner has derived this for pure silicon (70). His results can be easily extended to SiC and Si3N4 (71). The active-to-passive transition occurs when sufficient SiO(g), via Eq. (23), is generated to satisfy the SiC/SiO2 condition for equilibrium. There is some controversy about the exact equilibrium condition, but reasonable agreement with measurements is obtained from the following equilibria: 2SiC ⫹ SiO2 ⫽ 3Si ⫹ 2CO(g) (24) The transition P(O2) for active-to-passive oxidation is calculated based on equilibrium conditions and diffusion through the gas boundary layer to the sample. Calculated and measured active-to-passive transitions (72–74) are shown in Fig. 12. Now, consider the case of passive-to-active transition. Beginning with a stable SiO2 film on SiC; as the partial pressure of oxygen is lowered, the film Nonoxide Ceramics 331 Fig. 12 Calculated active-to-passive and passive-to-active transitions and measured active-to-passive transitions. (Data from Refs. 72–74.) will become unstable and SiO2 will react to form SiO(g). The transition pressure is calculated from the decomposition of the protective SiO2 film (70): SiO2(s) ⫽ SiO(g) ⫹ 1/2O2(g) (25) The calculated passive-to-active line is shown in Fig. 12—note that it is several orders of magnitude lower than the active-to-passive transition. A second route to SiO(g) formation occurs in a mixture of oxidizing and reducing gases (68). These can be present in a fuel-rich combustion situation 332 Jacobson and Opila with appreciable amounts of CO2, CO, H2O, and H2. It has also been observed in laboratory experiments with CO/CO2 (73) and H2 /H2O (72). Thus, a sequence of reactions might be SiC ⫹ 3H2O(g) ⫽ SiO2 ⫹ CO(g) ⫹ 3H2(g) (26) SiO2 ⫹ H2(g) ⫽ SiO(g) ⫹ H2O(g) (27) This leads to paralinear kinetics—simultaneous oxide growth and linear reduction kinetics. 2. Effects of H2S and Cl2 Hydrogen sulfide is found in coal gasifiers and various other petrochemical environments (5). At low oxygen potentials, when SiO2 does not form, both SiO(g) and SiS(g) may form: Si3N4 ⫹ 3H2O(g) ⫽ 3SiO(g) ⫹ 3H2(g) ⫹ 2N2(g) (28) Si3N4 ⫹ 3H2S(g) ⫽ 3SiS(g) ⫹ 3H2(g) ⫹ 2N2(g) (29) These reactions can lead to material consumption (5); however, when the oxygen potential is high enough to form SiO2, the corrosion rate drops substantially. Chlorine, which may be found in certain chemical process environments, also leads to volatile products (75). In the case of SiC, reaction occurs with the silicon, leading to formation of various silicon chlorides with residual carbon. As the oxygen potential is increased, attack becomes less severe (75). Current interest focuses on chlorine as HCl in waste incineration applications (76). Here, the situation is quite complex, involving not only HCl, but also a range of deposits. D. Oxidation in the Presence of Impurities and Deposits 1. Low Levels of Na and K Small amounts of impurities have several possible effects on the structure and, thus, the transport properties of silica. First, impurities can nucleate the formation of cristobalite at temperatures and times where amorphous silica would be expected. Second, impurities can modify the silica network, as shown in Fig. 3c. Breaking up the network tends to increase the oxygen transport rate through silica. Small amounts of alkali metals act as network modifiers and can increase oxidation rates an order of magnitude (56,77,78). 2. Na2SO4 Deposits As impurity levels increase, actual deposits may form. The chemistry of these depends on application, as shown in Table 1. Corrosive deposits have long been Nonoxide Ceramics 333 known to be an important issue for metals and alloys in various high-temperature applications (79). Here, we shall focus on sodium sulfate deposits on silica. Many of the general principles described for this system apply to the other cases. It is difficult to simulate the effect of an actual deposit. Simple laboratory experiments involve a one-time deposition of a salt on a test coupon, followed by a heat treatment. This type of experiment has the advantage of precise control of experimental variables such as temperature and gas composition. However, the actual situation involves a continuous deposition process. This can be done in a laboratory furnace. However, this can be accomplished more effectively by seeding a flame in a burner rig, as illustrated in Fig. 13. Figure 14 is a comparison of SiC treated in a burner without salt and with a flame seeded with NaCl. Note the extensive corrosion in the latter case. The formation of Na2SO4 occurs when ingested NaCl reacts with sulfur impurities in the fuel (80,81): 2NaCl(v) ⫹ SO3(g) ⫹ H2O(g) ⫽ Na2SO4(l) ⫹ 2HCl(g) (30) The source of NaCl depends on the application: It may be from a marine environment in the case of a heat engine, or from the process chemicals in the case of an industrial furnace. Under the appropriate conditions, Na2SO4 forms as a condensed phase, depositing on parts. Figure 15 shows the calculated dew points Fig. 13 Schematic of jet fuel burner used for corrosion studies. 334 Jacobson and Opila Fig. 14 Optical micrographs of sintered SiC coupons with carbon and boron additives, oxidized in a burner rig at 1000°C (a) 46 h with no sodium and (b) 13.5 h with a sodiumchloride-seeded flame. (From Ref. 24.) for Na2SO4 deposition (81). The rates of deposition are also critical and have been treated in detail (82). The most useful interpretation of this process is with the acid–base theory of molten salts. Na2SO4 decomposes to Na2O, which is the key reactant: Na2SO4 ⫽ Na2O ⫹ SO3(g) (31) Note that the overpressure of SO3 sets the chemical activity of Na2O. A high activity of Na2O is a basic molten salt; a low activity of Na2O is an acidic molten salt. Because SiO2 is an acidic oxide, it is readily attacked by a basic molten salt: 2SiO2(s) ⫹ Na2O ⫽ Na2O⋅2(SiO2)(l) (32) Nonoxide Ceramics 335 Fig. 15 Calculated dew points for Na2SO4 deposition. (From Ref. 24.) This describes the product observed in Fig. 14: A solid, protective SiO2 layer has been replaced by a liquid nonprotective sodium silicate layer. The liquid layer allows rapid diffusion of oxygen inward and carbon monoxide outward, leading to accelerated oxidation. The accelerated oxidation provides additional SiO2 for reaction (32) and can lead to rapid consumption of the SiC. The conditions for reaction (32) to proceed can be calculated from basic thermodynamics. We can assume unit activities for the SiO2 and Na2O⋅2(SiO2)(l) and, hence, the threshold Na2O activity [a(Na2O)] for SiO2 dissolution is given by a(Na2O) ⫽ exp 冢 冣 ⫺∆G RT (33) Here, ∆G is the free-energy change for reaction (32). At 900°C, the minimum a(Na2O) is 10⫺11, meaning any activity greater than this will lead to silicate formation. Because the activity of Na2O is set by P(SO3) in reaction (31), this means that a P(SO3) of 5 ⫻ 10⫺5 bar or less will lead to silicate formation. These predictions hold in the actual burner, where the P(SO3) is set by the sulfur content of the fuel. Figure 16 shows a series of quartz coupons treated in a burner with a low- (0.05%) and high- (0.5%) sulfur fuel. The low-sulfur fuel has a low P(SO3) (basic salt) and, thus, allows silicate formation, whereas, the high-sulfur fuel has a higher P(SO3) (acidic salt), which suppresses silicate formation. 336 Jacobson and Opila Fig. 16 Quartz coupons treated in burner rig for 1 h with a 2 ppm sodium (as NaCl) seeded flame: (a) No. 2 diesel fuel (0.5% sulfur), 5 h; (b) Jet A fuel (0.05% sulfur). (From Ref. 81.) Actual ceramic systems are more complex than the quartz coupons used in the preceding experiment. Ceramic systems are silica films on SiC and Si3N4 and may contain additives such as refractory oxides and carbon. We have seen that carbon, in silicon carbide and as an additive, leads to extensive salt corrosion. Using an electrochemical sensor, it can be shown that carbon tends to drive so- Nonoxide Ceramics 337 dium sulfate more basic. The exact chemical mechanism for this is not clear, but it may involve a Na2S intermediate (81). 3. Mixed Sulfates, Vanadates, Slags This general concept of acid–base reactions in molten-salt corrosion extends to other deposits as well. Ingested sea salt leads to MgSO4 and CaSO4 deposits as well as Na2SO4 deposits (83). This mixed sulfate has been described by moltensalt solution models (84) and leads to Mg and Ca silicates. Impure fuels lead to vanadate deposits. V2O5 is an acidic molten salt and, thus, does not react with SiO2. However, it is reported to accelerate Si3N4 oxidation, possibly due to the solubility of SiO2 in V2O5, which may be further enhanced by the presence of a Y2O3 in Si3N4 (85). A more complex case is that of a molten slag, which may contain up to 10 different oxides. Here, it is more difficult to define basicity; however, a useful index has been the weight percent ratio of basic to acid oxides (86). The basic and acidic oxides encountered are Basic oxides: Na2O, K2O, MgO, CaO, Fe2O3 Acidic oxides: SiO2, Al2O3, TiO2 Fig. 17 CVD SiC oxidized for 1 h at 1800°C. (Courtesy of D. Fox, NASA Glenn.) 338 Jacobson and Opila Fig. 18 (a) Total pressure [P(SiO) ⫹ P(CO)] generated from SiC/SiO2 interactions. (b) Total pressure [P(SiO) ⫹ P(N2)] generated from Si2N2O/SiO2 interactions. (From Ref. 24.) Nonoxide Ceramics 339 As in the Na2SO4 case, the basic slag tends to cause more rapid material degradation than the acidic slag. Under some conditions, metal silicides may form with a basic coal slag (4). E. High-Temperature Effects Silica melts at 1723°C. Transport rates are very high in a liquid scale and rapid oxidation occurs, as shown in Fig. 17. Some of the bubbles in this sample are from escape of CO(g), but some are also from the interaction of SiC and SiO2 at high temperatures: SiC(s) ⫹ 2SiO2(s) ⫽ 3SiO(g) ⫹ CO(g) (34) Figure 18a is a plot of total pressures generated by this reaction (87). Note that for carbon saturated SiC, the pressures are much higher, leading to a lower upper use temperature. An analogous situation exists for Si3N4, but based on the Si2N2O interaction with SiO2: Si2N2O(s) ⫹ SiO2(s) ⫽ 3SiO(g) ⫹ N2(g) (35) The total pressure is shown in Fig. 18b. IV. REFRACTORY OXIDE COATINGS One possible solution to the preceding issues of volatility and molten-salt corrosion is the use of refractory oxide coatings on SiC and Si3N4. Refractory oxides are generally more chemically inert and this approach offers the possibility of combining the benefits of both materials. Table 2 lists some coefficients of thermal expansion (CTEs) for SiC, Si3N4, and refractory oxides. There is a good Table 2 Coefficients of Thermal Expansion of SiC, Si3N4, and Several Refractory Oxides Material SiC Si3N4 Mullite (3Al2O3⋅2SiO2) Alumina (Al2O3) Partially stabilized zirconia (0.08 Y2O3⋅ZrO2) Source: Adapted from Ref. 89. CTE (106 K⫺1) 5.2 3.2 5.4 9 10 340 Jacobson and Opila Fig. 19 Micrograph of polished cross section of mullite on SiC. (From Ref. 88.) match between SiC and mullite. Mullite coatings on SiC were first developed at Solar Turbines (San Diego, CA) and further developed at NASA Glenn (88,89). The critical processing issues are surface roughening of the SiC for adherence and application of a fully crystalline mullite coating. A mullite coating on a SiC coupon is shown in Fig. 19. There are clear advantages with mullite coatings in Na2O-induced corrosion due to formation of high-melting sodium-aluminosilicates, as opposed to lower-melting sodium silicates. However, the chemical activity of silica in mullite is only about 0.4, so reactions which volatilize SiO2 can still occur readily. It may be possible to apply other refractory oxides to SiC. As Table 2 shows, the problem of CTEs must be considered. One approach is graded coatings from SiC to a refractory oxide such as alumina or zirconia. V. HYDROTHERMAL CORROSION OF SiC AND Si3N4 Hydrothermal corrosion involves the attack by water at high temperatures and high pressures. Typical conditions are shown in Table 1. These conditions are often near the critical point of water. Table 1 also shows some potential applications in which these environments may be encountered. Another promising application of SiC and Si3N4 is a vessel for oxidation of waste materials to form safe products using supercritical water (90). Nonoxide Ceramics 341 The hydrothermal corrosion of SiC may form silica by the following reactions (91,92): SiC ⫹ 2H2O(g) ⫽ SiO2 ⫹ CH4(g) (36) SiC ⫹ 4H2O(g) ⫽ SiO2 ⫹ CO2(g) ⫹ 4H2(g) (37) Given nearly equal amounts of SiC and water vapor at high temperatures, free carbon can form (92,93): SiC ⫹ 2H2O(g) ⫽ SiO2 ⫹ C ⫹ 2H2(g) (38) This provides a method for synthesis of carbon films. However, in most applications, there is excess water and SiO2 is formed as a protective film. The hydrothermal corrosion of Si3N4 is similar and occurs by Si3N4 ⫹ 6H2O(g) ⫽ 3SiO2 ⫹ 4NH3(g) (39) At temperatures above 430°C, NH3(g) decomposes to N2(g) and H2(g). In both cases, the silica is eventually dissolved by the high-pressure water as SiO2 ⫹ OH⫺ ⫽ HSiO3⫺ (40) In the case of Si3N4, this leads to attack of the grain-boundary phases and dislodging of the grains (94). Methodical studies have been done on the effect of additives in Si3N4 on its hydrothermal corrosion behavior (13,94). The morphology of attack has been correlated with the type of additive and its resistance to hydrothermal attack. Ceramics which form oxides other than silica behave differently under hydrothermal conditions. AlN forms AlOOH (95). Other ceramics such as TiC and ZrC, which do not form protective oxides at high temperatures, also do not form protective oxides under hydrothermal conditions (91). VI. LIQUID METALS Ceramics have long been used as containers for processing of liquid metals. More recent applications involve use of liquid metals as a heat-transfer medium. A wide range of nonoxide ceramics have been considered for these applications (12,96,97). Dissolution of the ceramic and capillary action are two important issues in liquid metal corrosion (12). Dissolution can be diffusion controlled or interface controlled. Local variations in dissolution due to phase and structural differences can lead to surface roughening. Capillary action leads to internal at- 342 Jacobson and Opila tack of porous ceramics. More importantly it can lead to substantial grain-boundary attack, depending on the relation of the solid–liquid interfacial energy and the grain-boundary energy (12). VII. EFFECT OF OXIDATION/CORROSION ON MECHANICAL PROPERTIES It is evident from current materials research that we cannot view one phenomenon such as oxidation or corrosion alone, but rather must view its effect on the entire system. Nearly all the applications listed in Table 1 utilize the ceramic in some type of load-bearing situation. A critical question that arises is how oxidation and corrosion alter the ability of the ceramic to bear a load (98–112). Again, the emphasis here is on SiC and Si3N4. Two recent reviews summarize the critical issues involved in corrosion– mechanical property interactions (98,99). Many of the studies in this area deal with the change in mechanical properties after corrosion. However, these processes often act together. The oxidation rate for various types of Si3N4 is increased by the application of either compressive or tensile stresses (100,101). It is well established that short-term oxidation tends to blunt cracks and may actually increase strength (103). Long-term isothermal oxidation (5000 h) for SiC containing B, C, and Si leads to strength increases, whereas the same treatment for SiC containing Al2O3 and WC leads to strength decreases (104). Little strength degradation of a variety of commercial SiC and Si3N4 materials was observed for 3500-h burner rig tests with 12-min cycles (105). Recent research (106) shows that the effects of water vapor on strength are complex; however, in some cases, the rapid oxidation rates lead to flaw healing. In the case of stress rupture for Si3N4, water vapor has little effect (107). Molten-salt corrosion (108–110) leads to extensive pitting and strength reduction. Figure 20 is a fracture origin due to a corrosion pit from molten-salt attack (108). This indicates the need for tougher, greater flaw-resistant ceramics. In the case of Si3N4, it has been shown that corrosion reactions often affect the grain-boundary phase. Various mechanical properties are often related to the stability of this phase. Compositional changes in this grain-boundary phase may lower the threshold stress intensity for crack growth and also increase the creep rate. Corrosion reactions at a growing crack tip are quite important as evidenced by moisture-assisted crack growth of ceramics (111). Henager and Jones (112) have shown that the presence of Na2SO4 increases the slow crack growth velocity by a factor of 2 over air at 1573K. Similar effects occur in the case of hydrothermal corrosion. Pitting and dislodging of grains in various types of Si3N4 leads to substantial strength reduc- Nonoxide Ceramics 343 Fig. 20 Fracture origin for sintered α-SiC reacted with Na2SO4 for 48 h at 1000°C. The area noted in (b) is enlarged in (c). (c) The preferential grain-boundary attack is observed at arrows. (From Ref. 108.) 344 Jacobson and Opila tions. Analogous to the high-temperature situation, changes in the grain-boundary phase of Si3N4 lead to changes in the mechanical behavior. Weakening and dissolution of this phase leads to substantial strength reductions. Similarly, liquid metal attack of grain boundaries will lead to strength reductions (96). VIII. CONCLUSIONS Nonoxide ceramics are a promising class of materials for a wide range of applications at high temperatures. Before application, their interactions with the hightemperature environment must be well understood. The silica-forming ceramics (SiC and Si3N4) exhibit the best oxidation behavior and the focus of this chapter is on them. However, some information is included on AlN and BN, which show promise for several specific applications. High-temperature oxidation is critical to most applications. The oxidation of SiC is rate limited by the diffusion of oxygen inward through the growing silica scale. The oxidation of Si3N4 is more complex: Here, the oxynitride scale plays a role. Most engineering ceramics contain additives, which tend to increase oxidation rates. Actual applications involve complex gas mixtures. The effects of water vapor, carbon dioxide, low oxidant pressures, and mixed oxidizing/reducing gases are discussed. Water vapor enhances oxidation, whereas carbon dioxide is a less effective oxidant. Low oxidant pressures and mixed oxidizing/reducing gases lead to SiO(g) formation. Corrosive deposits are also encountered in some applications. These include sodium sulfates, vanadate, and slags. Deposit-induced corrosion can be described with the acid–base theory of oxides. Because silica is an acidic oxide, it is readily attacked by a basic molten salt. These problems may be minimized with the application of a refractory oxide coating. Mullite shows a good thermal expansion match to SiC and is a promising starting point. It may be possible to develop other more refractory oxide coatings. Some proposed applications involve hydrothermal conditions: water at high temperatures and pressures. Hydrothermal corrosion is analogous to high-temperature oxidation in many ways. Ceramics which form protective high-temperature oxides also form protective hydrothermal oxides. The form of corrosive attack is also similar. Finally, the interaction of corrosive attack and degradation of mechanical properties is discussed briefly. 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Molten salt corrosion of hot-pressed Si3N4 /SiC-reinforced composites and the effects of molten salt exposure on slow crack growth of hot-pressed Si3N4. Ceram Trans 10:197–210, 1990. 11 Oxide Ceramics F. S. Pettit, G. H. Meier, and J. R. Blache`re University of Pittsburgh, Pittsburgh, Pennsylvania I. INTRODUCTION In discussing environmental effects on oxide ceramics, it is necessary to select a manageable number of environments and oxide systems to illustrate most of the salient points while also attempting to provide some systematic logic to the degradation effects in general. An oxide ceramic is a ceramic that contains a substantial amount of oxygen as a component. We can have binary oxides, ternary oxides, as well as more complex oxides. Examples of nonoxide ceramics are Si3N4 and SiC. In this chapter, we will consider environmental effects on binary oxides, namely Al2O3, Cr2O3, and SiO2. Environments affect oxide ceramics when the thermodynamic conditions are such that reactions occur between the environment and the ceramic. The extent of these reactions is determined by kinetic processes, as well as by the driving force for reaction to occur. Therefore, the particular way that the environmental conditions are imposed upon the ceramic play an important role. For example, the forms of the environmental effects are significantly different for oxides exposed to gas environments compared to liquid environments. In this chapter, we will consider environmental effects induced by gas environments—in particular, O2, N2, and some gas mixtures such as O2 ⫹ SO2, O2 ⫹ CO2, and H2O ⫹ O2. The liquid environments will be considered by using molten salts and liquid metals. As discussed by McCauley (1), oxide ceramics can also be affected by solid environments. High-temperature acid–base reactions between solids such as SiO2 and MgO are major problems which can lead to catastrophes by the reaction between refractories placed in contact. The low melting salts formed in the reaction can lead to the ‘‘melt down’’ of structures. 351 352 Pettit et al. Not placing acid and basic refractories in contact is a major requirement in the design of furnaces such as arc furnaces for steel-melting and glass-melting furnaces. Although reactions between solid oxide ceramics are important, such reactions will not be considered specifically in this chapter. However, the processes discussed in the liquid environment section are relevant to solid reactions. Finally, it is important to emphasize that environmental effects on ceramics can extend from room temperature to extremely high temperatures. In this chapter, a very wide range of temperatures will be used to illustrate the various environmental effects. II. GASEOUS ENVIRONMENTS A. Oxygen Most oxide ceramics are rather insensitive to oxygen, especially at low temperatures. As the oxygen pressure is varied, oxide ceramics will adjust their stoichiometry or transform to higher or lower oxide phases at rates commensurate with the temperature and the oxygen pressure differential. In the case of changes in stoichiometry, even though the changes may be small, the effects on properties may be substantial. This is shown in Fig. 1A, where α-Al2O3 can be either an n- or p-type electronic conductor, or an ionic conductor, depending on the oxygen pressure and temperature (2). A Kro¨ger–Vink diagram showing defect concentrations that could account for the observed conductivities at 1400°C is also included in Fig. 1B. Comparison of Figs. 1A and 1B show that at 1400°C for oxygen pressures of 1 atm, the principal point defects are cation vacancies and electron holes, and Al2O3 is a p-type electronic conductor. As the oxygen pressure is decreased, the principal defects become aluminum ion interstitials and cation vacancies and the Al2O3 is an ionic conductor. At even lower oxygen pressures, Al2O3 becomes an n-type electronic conductor. For some oxide ceramics, as the oxygen pressure is changed, another oxide phase can be formed. For example, Cu2O could be converted to CuO. To determine if such phase changes may occur, standard free energies of formation can be used to calculate the equilibrium pressures for the phases of concern. For example, in the case of Cu2O and CuO, the equilibrium oxygen pressure can be determined as follows: Cu2O ⫹ 1/2O2 → 2CuO P O2 ⫽ exp 冢 (1) 冣 2∆G °CuO ⫺ ∆G °Cu2O 2RT P O2 ⫽ 10⫺4 atm at 1000 K (2) Oxide Ceramics 353 Fig. 1 Diagram showing the type of conduction predominating in α-Al2O3 as a function of temperature and oxygen pressure (A), and a Kro¨ger–Vink diagram to account for these conductivity regimes (B). 354 Pettit et al. The effect of the oxygen pressure change depends on the function the oxide ceramic may be serving. In the case of Cr2O3, which is often used to develop oxidation resistance on nickel-, cobalt-, and iron-based alloys, the following reaction must be considered: Cr2O3 ⫹ 3/2O2 → 2CrO3(g) (3) which gives PCrO3 ⫽ P O3/42 exp 冢 冣 ∆G °Cr2O3 ⫺ 2∆G °CrO3 2RT (4) Equation (4) shows that the pressure P CrO3 increases with oxygen pressure. Therefore, the formation of CrO3 causes Cr2O3 to be a less effective protective oxide barrier and this condition increased with oxygen pressure and temperature increases (∆G °Cr2O3 ⫺ 2∆G °CrO3 is negative). The vapor pressures of oxide phases are always a factor that must be considered. The vapor pressure of Al2O3 is relatively low and the pressures of other oxides involving aluminum and oxygen are also low, and so, Al2O3 is an oxide ceramic that can be used at temperatures up to 1300–1400°C with little effects of vaporization. On the other hand, MoO3 has high vapor pressures and cannot be used at temperatures above 400–500°C. Gulbransen and Jansen (3) have proposed that vaporization of oxides begins to cause problems at pressures greater than 10⫺9 atm. An interesting situation arises in the case of SiO2, as has been discussed by Wagner (4). As indicated by the equations SiO2 → SiO(g) ⫹ 1/2O2 P SiO ⫽ exp [(∆G °SiO2 ⫺ ∆G °SiO)/RT] P O1/22 (5) (6) the pressure of SiO(g) increases as the oxygen pressure decreases. Therefore, at low oxygen pressure, SiO formation can affect the oxidation of silicon, as well as some nonoxide ceramics such as Si3N4. When silica glass is exposed to elevated temperatures, it can devitrify or change from the glass state to the crystalline state (5). Devitrification results in significant changes in the properties of the silica. The cause of devitrification is not completely understood. High temperatures are certainly a factor, but oxygen in the gas, or water vapor, have been proposed to be also necessary. B. Pure Gas Environments Other Than Oxygen When oxide ceramics are exposed to gases other than oxygen, depending on the conditions, the oxide ceramics can be changed to phases determined by the other Oxide Ceramics 355 gas. For example, when the gas is pure nitrogen and the oxide ceramic is Al2O3, the following reaction can be used to determine what will occur: Al2O3 ⫹ N2 → 2AlN ⫹ 3/2O2 PN2 P O3/22 ⫽ exp 冢 (7) 冣 2∆G °AlN ⫺ ∆G °Al2O3 RT (8) At 1500 K, PN2 /PO3/22 ⫽ 3 ⫻ 1030, and, therefore, if the P N2 /P O3/22 ratio in the gas is greater than this value, the Al2O3 will be converted to AlN. Of course, the rate at which this will occur will be determined by kinetics, which, in turn, is determined by factors that usually require some experimentation to understand. For example, the AlN will form as a layer covering the Al2O3 and the reaction will involve nitrogen diffusing inward and oxygen diffusing outward. If the oxygen cannot be removed, pressures could develop that could rupture the AlN layer and cause the rate of the reaction to be changed. In principal, reactions, such as those described for nitrogen with Al2O3, could occur for sulfur and carbon, as well as numerous other gases, but it is important to emphasize that almost all metals have much greater affinities for oxygen than for these other reactants, and, therefore, the oxide ceramics are comparatively stable. At 1500 K, the ratio PN2 /P O3/22 for equilibrium between Al2O3 and AlN is 3 ⫻ 1030. At a total pressure of 1 atm, if the nitrogen gas contains even a small amount of oxygen, the oxide phase will remain stable. C. Mixed Gases Containing Oxygen When oxide ceramics are exposed to gases containing oxygen and another reactant, as discussed previously, the oxygen usually is not displaced unless the oxygen pressure is extremely low. Even though the oxygen is not displaced from the ceramic, the second reactant can affect the oxide ceramic by reactions of the following type, where sulfur is the second gaseous reactant: Al2O3 ⫹ 9/2O2 ⫹ 3/2S2 → Al2O3 ⫹ 3SO3 → Al2(SO4)3 (9) The conditions for which such reactions can occur can be described by constructing thermodynamic diagrams such as the one presented in Fig. 2 for the aluminum–oxygen–sulfur system. Again, the rate of conversion of Al2O3 to Al2(SO4)3 (Fig. 2) will be determined by kinetic processes. Reactions such as that described by Eq. (9) can occur for carbon and nitrogen, as well as other gases. Usually, most gas environments encountered in industrial practice are not appropriate for reactions of this type to take place. However, the stabilities of these phases do vary. An interesting example is CaO or MgO, which is often present in Al2O3 as impurities. The sulfates of these two oxides are more stable than 356 Pettit et al. Fig. 2 Stability diagram for Al2O3 as a function of O2 and SO3 pressures at 700°C where the boundaries for the CaO/CaSO4 and MgO/MgSO4 are also included. Al2(SO4)3, and these sulfates have been observed (5) to form on Al2O3 without any Al2(SO4)3 formation (Fig. 3). In some gas mixtures, the second reactant may not enter directly into a chemical reaction with the oxide ceramic, but it still exerts a profound effect on the mechanical properties of the ceramic. An excellent example is the effect of water vapor on crack propagation in soda-lime silica glass as has been described in articles by Weiderhorn (6,7). The strength of glass is known to deteriorate in atmospheres containing moisture. Under constant tensile load below the fracture strength at room temperature, the glass may eventually fail in a process called static fatigue. This delayed fracture is associated with preferential attack by water of stressed SiEO bonds at the crack tip. This is usually considered a stress corrosion associated with the reaction between the glass and water, as indicated schematically in Fig. 4, in which an oxygen bridge of the glass structure is broken. This was proposed by Orowan and many others. Wiederhorn found that the velocity of crack propagation in soda-lime silica glass depended on the humidity according to three regimes as shown in Fig. 5. Over a range of stresses, the velocity of crack propagation depended on the relative humidity as expected from Fig. 4. However, at higher stresses, the velocity of the crack propagation was limited apparently by the rate of transport of the water to the crack tip. This transport was usually assumed to be by surface diffu- Oxide Ceramics 357 Fig. 3 Scanning micrograph of alumina exposed to 7 ⫻ 10⫺3 atm of SO3 in O2 at 700°C where Mg(SO4) and Ca(SO4) were detected but not Al2(SO4)3. Fig. 4 Schematic representation of proposed reaction between water and a strained SiEOESi bond at a crack tip. Step 1 involves adsorption of water to the SiEO bond and step 2 shows the reaction involving simultaneous proton and electron transfer. In step 3, surface hydroxyl groups are formed. 358 Pettit et al. Fig. 5 Dependence of crack velocity in silica on the applied force. The percent relative humidity for each set of runs is given on the right-hand side. Roman numerals identify the different regions of crack propagation. In region I, the crack growth rate is proposed to be controlled by adsorption or reaction of water vapor at the crack tip. In region II, the rate is supposedly controlled by diffusion of water vapor in the crack. Region III involves a crack propagation mode not involving water vapor. sion along the crack or by vapor transport. At higher stresses yet, the crack velocity does not depend on humidity. Tomazoa (8) recently discussed that either crack initiation or crack growth can dominate static fatigue. He emphasized the evidence for water penetration into the glass. The water would diffuse into the glass as molecular water and react with the glass structure as shown in Fig. 4. The interaction of water with the strength of glass is well established and is expected to affect the strength of many ceramics which contain a glassy phase. Michalski and Freiman (9) proposed that such processes can take place in other oxides and species containing polar molecules. It has been found that such an Oxide Ceramics 359 effect can occur in Al2O3. Furthermore, similar processes have been proposed as the cause of the excessive spalling of alumina scales from metallic alloys exposed to gases containing water vapor (10). III. LIQUID ENVIRONMENTS In discussing the effects of liquid environments on oxide ceramics, a variety of types of liquid must be considered, and the thickness of the liquid is also an important factor affecting the types of interaction that can occur. For example, in the case of thin liquid molten-salt deposits, the gas environment above the liquid deposit can enter into the reaction scheme. In this chapter, the effects of thin liquid molten-salt deposits will be considered first and then additional effects arising from thin deposits of other types of liquids will be discussed. Finally, the changes in these types of interactions as the liquid becomes very thick will be examined. A. Thin Liquid Deposits of Molten Salts Hot corrosion (11–13) is a process whereby liquid deposits, usually molten deposits such as sulfates and vanadates, cause corrosion of metals and ceramics. In the case of the hot corrosion of oxide ceramics, three types of process must be considered (14): • Reactions between the ceramic and the deposit • Diffusion of various constituents through the molten deposit • Reactions between the gas and the molten deposit These three processes are coupled to each other, with the slowest determining the overall reaction rate. To comment on specific processes that may be occurring, it is necessary to consider some specific oxides, liquid deposits, and gas compositions; but before doing that, some generalizations are worthwhile. 1. Theoretical Considerations The most important reactions between oxide ceramics and molten deposits can be viewed as involving oxide ions, where considering the metal M, oxide ions can be donated to the melt, MO2 → M4⫹ ⫹ 202⫺ (10) or can be taken from the melt MO2 ⫹ O2⫺ → MO32⫺ (11) 360 Pettit et al. Depending on the characteristics of the melt, the oxide ions may not exist as major species but can be defined by reactions of the type SO3 ⫹ O2⫺ → SO42⫺ (12) 2SO3 ⫹ O2⫺ → S2O72⫺ (13) V2O5 ⫹ O2⫺ → 2VO3⫺ (14) In view of such reactions, at times it is useful to view the oxide ions as a basic component in the melt where in the previous equations the acidic components would be SO3 or V2O5. Stability diagrams can be used to identify the reactions that may occur between an oxide ceramic and a molten deposit. For example, in the case of a Na2SO4 deposit, the diagram in Fig. 6 defines the stability range of Na2SO4 at a fixed temperature in terms of oxygen and SO3 pressures. The SO3 pressure is the acidic component in the melt and is related to the basic component, Na2O, or O2⫺, via the equation Na2O ⫹ SO3 → Na2SO4 (15) Fig. 6 Stability diagram for the Na–O–S system at 1000°C; the dashed lines are sulfur isobars. Oxide Ceramics 361 where for pure Na2SO4 K 15 ⫽ 1 a Na2OP SO3 (16) The basic component can also be considered to be oxide ions because SO42⫺ → SO3 ⫹ O2⫺ (17) The stability diagram for the oxide ceramic can now be superimposed upon the stability diagram for Na2SO4. In Fig. 7, the stability diagram for the Al–O–S system has been superimposed on Fig. 6. Inspection of Fig. 7 shows that depending on the composition of the Na2SO4, aluminum may be present as NaAlO2 (AlO2⫺ ions), Al2O3, Al2(SO4)3 (Al3⫹ ions), or Al2S3. It is important to note that these are the phases of aluminum which exist at unit activity. Hence, in the region of Fig. 7 where Al2O3 is stable, an important equilibrium is Al2O3 → 2Al3⫹ ⫹ 3O2⫺ (18) Fig. 7 Stability diagram showing the phases of aluminum that are stable in Na2SO4 at 1000°C; the dashed lines indicate sulfur isobars and the arrows indicate compositional changes of Na2SO4. 362 Pettit et al. where K 18 ⫽ a 2Al3⫹a 3O2⫺ and, therefore, as the activity of oxide ions decreases, the activity of Al3⫹ ions increases until a value of unity is reached at the Al2O3 –Al3⫹ boundary. Another reaction of importance is Al2O3 ⫹ O2⫺ → 2AlO2⫺ (19) and the activity of AlO2⫺ ions increases as the oxide ion activity increases until a value of unity is reached at the AlO2⫺ –Al2O3 boundary. Rapp (13) has experimentally determined the solubilities of a number of oxides in Na2SO4, and some of the results from these studies are presented in Fig. 8. The slopes of these solubility curves can be used to verify the reactions by which the oxide dissolves. For example, in the case of Al2O3, basic reactions can be written as Al2O3 ⫹ Na2SO4 → 2NaAlO2 ⫹ SO3 (20) Al2O3 ⫹ O2⫺ → 2AlO2⫺ (21) or which both yield slopes of ⫺1/2. The acidic dissolution reactions can be written as Al2O3 ⫹ 3Na2SO4 → Al2(SO4)3 ⫹ 3Na2O (22) Fig. 8 Solubilities of some oxides in Na2SO4 at 1200 K as determined by Rapp (13). Dissolution reactions consistent with the solubility curves for Al2O3 are included on this diagram. The dashed lines give calculated solubilities for Al2O3 obtained by using Temkin’s rule for ionic melts and the indicated reactions for Al2O3. Oxide Ceramics 363 or Al2O3 → 2Al3⫹ ⫹ 3O2⫺ (23) which both have slopes of ⫹3/2. When considering more complex melts, such as solutions of Na2SO4 – NaVO3, it is necessary to relate the two acidic components, SO3 and V2O5. This can be done by using equations such as Na2SO4 ⫹ V2O5 → 2NaVO3 ⫹ SO3 (24) where K 24 ⫽ (1 ⫺ X Na2SO4)2PSO3 X Na2SO4a V2O3 (25) and K 24 is the equilibrium constant, XNa2SO4 is the mole fraction of Na2SO4 in the Na2SO4 –NaVO3 solution, which is assumed to be ideal, PSO3 is the pressure of SO3 and aV2O5 is the activity of V2O5 in equilibrium with the melt. In order to use Eqs. (24) and (25), it is necessary to determine the major components in the melts of interest. For example, the vanadium may exist as V2O5, NaVO3, Na4V2O7, or Na3VO4. Equilibrium conditions for reactions such as V2O5 ⫹ 4Na3VO4 → 3Na4V2O7 (26) V2O5 ⫹ Na4V2O7 → 4NaVO3 (27) 2NaVO3 → Na2O ⫹ V2O5 (28) can be used to determine the proportions of the important species. Rapp (13) has developed diagrams for solutions of Na2SO4 –NaVO3 and typical results are presented in Fig. 9 for a temperature of 900°C. These results show that for SO3 pressures between 10⫺2 and 10⫺7 at 900°C, Na2SO4 –NaVO3 melts contain predominantly NaVO3. The amount of NaVO3 in the melt versus Na3VO4, Na3V2O7, or V2O5 also depends on temperature. When Na2SO4 and NaVO3 are the major species in Na2SO4 –NaVO3 melts, Eq. (25) can be used to relate the activities of the two acidic species, SO3 and V2O5. The solubilities of oxides in Na2SO4 –NaVO3 melts are influenced by both the basic and acidic components. In Fig. 10, the solubilities of some oxides in Na2SO4 and in a Na2SO4 –30 mol% NaVO3 melt are compared using results obtained by Hwang and Rapp (15). Because the basic component is the same in both melts, the solubility curves are identical for both melts in regions where dissolution involves the basic component. The differences arise when the dissolution reactions involve acidic species with higher solubilities occurring for V2O5 than SO3. 364 Pettit et al. Fig. 9 Mole fractions of V2O5, NaVO3, Na2V2O7, and NaVO3 in a Na2SO4 –20 mol% NaVO3 melt at 900°C as a function of SO3 pressure. Another point to be considered in molten deposits on alloys and ceramics involves the ionic species that are responsible for transport. In the case of Na2SO4 deposits, Na⫹ and SO42⫺ ions predominate, but it has been shown (11) that the transport of SO3 occurs via S2O72⫺ ions. In Na2SO4 –NaVO3 melts, the preceding ions are also important, but in view of the results presented in Fig. 9, VO3⫺, V2O73⫺, and VO43⫺ ions are important, with VO⫺ 3 ions having larger concentrations than V2O74⫺ or VO43⫺ in the melts exposed to SO3 pressures in the range between 10⫺2 and 10⫺7. Two final points worth noting in the solubility curves of oxides in molten deposits are dissolution–reprecipitation processes (16) and synergistic dissolution (17). Dissolution–reprecipitation occurs when the gradient of the oxide solubility is negative at the oxide–melt interface. This means that the oxide solubility is Oxide Ceramics 365 Fig. 10 Solubilities of Al2O3 in pure Na2SO4 and in Na2SO4 –30 mol% NaVO3 at 900°C as determined by Hwang and Rapp (15). greatest at the oxide–melt interface. The oxide can, therefore, dissolve but precipitate out in the melt where the solubility is lower. Synergistic dissolution occurs when two oxides react with melts in ways that the dissolution of one accelerates the dissolution of the other. For example, if the solubility curves for Al2O3 and Fe2O3 in Fig. 8 are considered at a value of log aNa2O ⫽ ⫺14, the dissolution of Al2O3 occurs via Al2O3 ⫹ O2⫺ → 2AlO2⫺ (29) As this reaction proceeds, oxide ions are consumed, and the melt becomes more acidic. The Fe2O3 dissolution involves the reaction / Fe2O3 → 2/3Fe3⫹ ⫹ O2⫺ 13 (30) where oxide ions are donated. The sum of these two reactions is /3Fe2O3 ⫹ Al2O3 → 2AlO2⫺ ⫹ 2/3Fe3⫹ 1 (31) The combined dissolution process of both oxides is such that the melt composition does not change with respect to oxide ion concentration and, hence, both reactions proceed more rapidly when occurring concomitantly. 2. Alumina Corrosion The corrosion of Al2O3 when exposed to deposits of molten salts is dependent on purity (18,19). High-purity (99.9% pure) α-Al2O3 single crystals lost weight 366 Pettit et al. as a function of time when exposed to deposits of Na2SO4 and Na2SO4 –NaVO3 in gas mixtures of oxygen and SO3 at temperatures of 700°C and 900°C. In the sulfate melts, the weight-loss rates decreased as the SO3 pressure was increased, which shows that the dissolution process is probably Al2O3 ⫹ SO42⫺ → 2AlO2⫺ ⫹ SO3 (32) The alumina dissolves by basic dissolution. The SO3 pressure is greater at the Al2O3 interface than the gas interface and SO3 is lost from the deposit. As dissolution of the alumina continues, the increased concentration of AlO2⫺ ions causes the SO3 at the Al2O3 interface to approach that of the gas interface, and, eventually, the dissolution process will stop unless fresh Na2SO4 is applied. The weight-loss rates of Al2O3 increased as the activity of V2O5 was increased in Na2SO4 –NaVO3 deposits. The dissolution reaction is proposed to be Al2O3 ⫹ 3VO3⫺ → 2Al3⫹ ⫹ 3VO43⫺ (33) with the following reaction occurring at the melt–gas interface: VO43⫺ ⫹ SO3 → VO3⫺ ⫹ SO42⫺ (34) This observed dissolution of Al2O3 in the Na2SO4 –NaVO3 melt corresponds to melt compositions on the acidic dissolution side of the solubility curve in Fig. 10, with the melt at the gas interface being more acidic than the melt at the Al2O3 interface. Impurities can play a significant role in the hot corrosion of Al2O3. A variety of conditions can occur depending on the type of impurity, the concentration, the melt composition, the gas composition, and the temperature. In the case of polycrystalline α-Al2O3 of 99.8% purity which contained SiO2, MgO, CaO and Na2O as impurities, a hot corrosion attack caused a porous zone to be developed (Fig. 11). The thicknesses of these zones conformed to the parabolic rate law and rate constants were dependent on the activity of V2O5 in the molten deposits (Fig. 12) with little effect of temperature. These polycrystalline specimens were attacked more rapidly than the high-purity single crystals due to the impurities present in the former specimens. It is also important to note that for activities of V2O5 above 10⫺5, transport through the melt via VO3⫺ ions determines the rates. However, at lower values, the rates are controlled by transport via S2O72⫺ ions. A good example of impurity effects has been observed (18) during the hot corrosion of Al2O3 induced by Na2SO4 in air at 700°C and 1000°C. Under these conditions, no attack of the high-purity single crystals was observed. ‘‘Highpurity’’ polycrystalline Al2O3 (0.1% MgO, 0.1% SiO2), however, developed features showing that some attack had occurred. At 700°C, some preferential dissolution of the grains was evident and some silicon-rich needles had developed on the surface of the Al2O3 (Fig. 13a). At 1000°C, a consistent pattern of sodium aluminum silicate along grain boundaries and sodium magnesium aluminum sili- Oxide Ceramics 367 Fig. 11 Cross section of polycrystalline alumina specimen exposed to hot corrosion conditions at 700°C that caused a porous zone to be developed. cate at triple points was evident (Fig. 13b). These results show that the Na2O activity in these melts is high enough to promote reaction with silicates present at grain boundaries of the Al2O3. As Na2O is removed from the deposit, the SO3 pressure is increased to levels at which Al2O3 can dissolve by acidic dissolution. Synergistic dissolution has occurred. 3. Silica Corrosion The solubility for vitreous silica in acidic melts is extremely low (20), but silica does react with basic melts to form Na2SiO3 (21) and there are significant solubilities (22). Furthermore, the silica devitrifies under certain conditions and sodium in melts accelerates this devitrification (23). In view of these conditions, silica is very resistant to acidic melts but rather susceptible to hot corrosion by basic melts. Typical reaction products developed on silica with a Na2SO4 deposit in air for different times at 1000°C are presented in Figs. 14 and 15. After 1 h, the Na2SO4 wetted the silica and a layer of cristobalite spherulites formed under the 368 Pettit et al. Fig. 12 Parabolic rate constants for growth of porous zones on polycrystalline alumina as a function of activity of V2O5 in the deposit. Horizontal lines indicate the rate constants for sulfate deposits. salt. After 24 h, nothing remained of the spherulitic surface morphology because tridymite formed at the silica–salt interface. Cristobalite separated the tridymite from the vitreous silica. After 10 h, a thin sodium silicate layer over tridymite and cristobalite layers was evident (Fig. 15). The crystalline layers spalled extensively on cooling from the test temperature as anticipated from their phase transitions. The major reactions between Na2SO4 deposits and vitreous silica consist of formation and dissolution of silicate via SiO2 ⫹ Na2O → Na2SiO3 (35) and diffusion of sodium into vitreous silica. In some deposits, the SiO32⫺ ions can convert to silica out in the melt when the acidity is sufficiently high. Sodium does diffuse into fused silica and may be incorporated in the glass as a network modifier (24) via the reaction Na2O ⫹ SiEOESi → 2SiEO⫺ ⫹ 2Na⫹ (36) The penetration of Na2O is expected to increase rapidly with the activity of Na2O in the melt because the driving force is increased. Initially, the sodium must enter the glass via reaction (36), which may be rate controlling, but the situation is modified by crystallization of the silica glass. Diffusion of sodium through cristobalite is very difficult, because it does not contain the channels of vitreous silica. Oxide Ceramics 369 Fig. 13 High-purity polycrystalline alumina exposed in basic conditions at (a) 700°C for 24 h and (b) at 1000°C for 405 h. (a) Silica-rich needles and (b) silicate reaction products form at grain boundaries on washed substrates. 370 Pettit et al. Fig. 14 Fused silica exposed to basic conditions at 1000°C for 1 h developed globular arrays of spherulites. B. Molten Metals The attack of oxide ceramics by molten metals has been discussed by McCauley (1). This type of attack is encountered in the manufacture of steel and nonferrous metals such as aluminum and copper. A primary cause of the corrosion process is reaction of the molten metals with oxygen in the oxide ceramic. Consequently, the thermodynamic stability of the oxide ceramic compared to that of the metal oxide is an important factor. In the case of alumina exposed to some metal M, the important reaction is Al2O3 ⫹ M (molten metal) → 3MO ⫹ 2Al (molten metal) (37) with the equilibrium conditions 冢 冣 a 2Al aM eq ⫽ exp [⫺(3∆G °MO ⫺ ∆G °Al2O3)] RT (38) where it is assumed that there is no solubility of MO in Al2O3 and that no ternary oxides (e.g., MAl2O4) are formed. If the standard free energy of MO per gramatom of oxygen is large compared to that for alumina, the right-hand side of Eq. Oxide Ceramics 371 Fig. 15 Cross section of fused silica exposed to basic conditions at 1000°C for 100 h with thermal cycling; three layers cover the smooth glass: sodium silicates, tridymite, and cristobalite, which is in contact with the glass. (38) can be a large positive number and so conditions will be favorable for Eq. (37) to proceed to the right. Even when MO is less stable than alumina, in principle some aluminum must be in the molten metal, and so some reaction can occur. The extent of oxide ceramic–metal reactions is determined by the distribution of the reaction product MO. When MO forms as a continuous layer over the alumina, the reaction rate will be controlled by transport through MO and can be expected to be small. However, if the MO is discontinuous or as particles in the molten metal, then reaction rates can be rapid. C. Corrosion in Thick Melts Corrosion of oxide ceramics in deep melts involves many of the same reactions that have been discussed previously, but the gas phase usually does not play a significant role. In cases where a component in the gas phase is involved in the corrosion process, this component can be substantially depleted from the liquid phase, and, consequently, the corrosive characteristics of the liquid can change substantially. Such conditions usually are established during the corrosion of metals and alloys in thick melts where, for example, oxygen is depleted from molten salts. To illustrate the effects of corrosion of oxide ceramics, only examples where the gas phase does not play a significant role will be considered. Many of the 372 Pettit et al. Fig. 16 Schematic to show the products developed during exposure of sapphire in CaO– MgO–Al2O3 –SiO2 slags. factors that affect the corrosion or dissolution of oxides in liquids have been described by Cooper and Kingery (25) in studies concerned with the dissolution of sapphire in CaO–Al2O3 –SiO2 slags, where it was shown that rates were controlled by mass transport in the liquid. The dissolution of sapphire in this slag was controlled by mass transport in the liquid even when transport rates were increased by rotating the immersed specimens. These investigators also analyzed the effects of molecular diffusion, natural convection, and forced convection with self-consistent results. Sandage and Yurek (26,27) studied the corrosion of sapphire and (Al, Cr)2O3 in CaO–MgO–Al2O3 –SiO2 slags and showed that a solid spinel reaction product (MgAl2O4) could be formed upon the surface of the sapphire depending on the experimental conditions. At times, the spinel did not form a continuous layer, but it did affect the dissolution of the sapphire. It was shown that in cases where the spinel formed as a continuous layer upon the sapphire, a steady-state condition was achieved for which a constant spinel thickness was established. The proposed mechanism for this process is shown schematically in Fig. 16. The net effect of this process is to transfer Al2O3 from the sapphire to the melt, but because the dissolution process involves transport through an intermediate phase, it is called indirect dissolution. IV. CONCLUDING REMARKS The effects of various environmental conditions on oxide ceramics are not remarkably different from nonoxide ceramic or metallic alloys. The differences are in the magnitude of the effects, due to the stability of oxide ceramics compared to nonoxide ceramics and especially metallic alloys in the environments that are usually encountered in practice. Oxide Ceramics 373 REFERENCES 1. RA McCauly. Corrosion of Ceramics. New York: Marcel Dekker, 1995, pp. 194– 197. 2. K Kitazawa, RL Coble. Electrical conduction in single-crystal and polycrystalline Al2O3 at high temperatures. J Am Ceram Soc 57:245–250, 1974. 3. EA Gulbransen, SA Jansen. Thermochemistry of gas–matal reactions. In: DL Douglass, ed. Oxidation of Metals. Materials Park, OH: American Society for Metals, 1971, pp. 63–86. 4. C Wagner. Passivity during the oxidation of silicon at elevated temperatures. J Appl Phys 29:1295, 1958. 5. HR Kim. Gaseous hot corrosion of oxide ceramics. MS dissertation, University of Pittsburgh, Pittsburgh, PA, 1983. 6. SM Wiederhorn. Moisture assisted crack growth in ceramics. Fracture Mech 4:171– 177, 1968. 7. SM Wiederhorn. Influence of water vapor on crack propagation in soda-lime glass. J Am Ceram Soc 50:407–444, 1967. 8. M Tomozawa. Fracture of glasses. Annu Rev Mater Sci 12:43–74, 1996. 9. TA Michalski, SW Freiman. A molecular mechanism for stress corrosion in vitreous silica. J Am Ceram Soc 66:284–288, 1983. 10. RJ Janakiraman, GH Meier, FS Pettit. The effect of water vapor on the oxidation of alloys that develop alumina scales for protection. Met and Mat Trans 30A:2905– 2913, 1999. 11. KL Luthra. Low temperature hot corrosion of cobalt-base alloys: Part I. Morphology of the reaction product. Part II. Reaction mechanism. Met Trans A 13A:1843–1864, 1982. 12. FS Pettit, CS Giggins. Hot corrosion. In: C Sims, N Stoloff, W Hagel, eds. Superalloys II. New York: John Wiley & Sons, 1987, pp. 327–358. 13. RA Rapp. Hot corrosion of materials. In: O Johannesen, A Andersen, eds. Selected Topics in High Temperature Chemistry. New York: Elsevier, 1989, pp. 291–329. 14. FS Pettit, GH Meier, JR Blachere. Hot corrosion of oxide ceramics. In: KG Nickel, ed. Corrosion of Advanced Ceramics. Dordrecht, The Netherlands: Kluwer, 1994, pp. 235–248. 15. YS Hwang, RA Rapp. Thermochemistry and solubilities of oxides in sodium sulfate–vanadate solutions. Corrosion 45:933–937, 1989. 16. RA Rapp, KS Goto. The hot corrosion of metals by molten salts. In: J Braunstein, JR Selman, eds. Proceedings of the Electrochemical Society Symposium on Molten Salts. The Electrochemical Society, Pennington, NJ, 1981, pp. 159–177. 17. YS Hwang, RA Rapp. Synergistic dissolution of oxides in molten salts. In: T Grobstein, J Doychak, eds. Oxidation of High Temperature Intermetallics. Warrendale, PA: TMS, pp. 257–270. 18. MG Lawson, FS Pettit, JR Blachere. Hot corrosion of alumina. J Mater Res 8:1964– 1971, 1993. 19. BM Warnes. The influences of vanadium on the sulfate induced hot corrosion of thermal barrier coating materials. PhD dissertation, University of Pittsburgh, Pittsburgh, PA, 1990. 374 Pettit et al. 20. DZ Shi, RA Rapp. J Electrochem Soc 133:849, 1986. 21. NS Jacobson. Sodium sulfate: Deposition and dissolution of silica. Oxide Metals 31:91–103, 1989. 22. GM Kim. The effects of contaminants and solubilities of SiO2 in fused Na2SO4 at 1200 K, MS dissertation, University of Pittsburgh, Pittsburgh, PA, 1982. 23. MG Lawson, HR Kim, FS Pettit, JR Blachere. Hot corrosion of silica. J Am Ceram Soc 73:989–995, 1990. 24. WD Kingery, HK Bowen, DR Uhlman. Introduction to Ceramics, 2nd ed. New York: John Wiley & Son, 1976, p. 103. 25. AR Cooper Jr, WD Kingery. Dissolution in ceramic systems: I. Molecular diffusion, natural convection, and forced convection studies of sapphire dissolution in calcium aluminum silicate. J Am Ceram Soc 47:37–43, 1964. 26. KH Sandage, GJ Yurek. Indirect dissolution of sapphire into silicate melts. J Am Ceram Soc 71:476–489, 1988. 27. KH Sandage, GJ Yurek. Indirect dissolution of (Al, Cr)2O3 in CaO–MgO–Al2O3 – SiO2 (CMAS) melts. J Am Ceram Soc 74:1941–1954, 1991. 12 Metal Matrix Composites Russell H. Jones Pacific Northwest National Laboratory, Richland, Washington I. INTRODUCTION Composite reinforcements are added to metals to increase their strength, elastic modulus, wear performance, thermal conductivity or to alter their thermal expansion properties. Applications for metal matrix composites (MMCs) include aerospace, automotive, military hardware, and so forth. The high specific properties are an obvious advantage to aerospace, but they are also very attractive for helping the automotive industry achieve its goals to build lighter-weight automobiles. Antenna, aircraft support structures, vertical tail fins, inertial guidance, and precision optical systems are some of the aerospace applications for MMCs as well as the cargo bay stiffeners on the space shuttle. Both continuous and discontinuous reinforcements are added to MMCs. Continuous reinforcements are usually graphite or boron fibers, whereas discontinuous reinforcements are usually SiC or Al2O3 particles or, in a few cases, whiskers. Graphite and boron fibers can induce galvanic corrosion, whereas SiC and Al2O3 particles are nonconducting and therefore will not induce galvanic corrosion. Clusters of particles can induce pitting corrosion, whereas both fibers and particles can induce localized corrosion. There is a considerable amount of corrosion data on aluminum matrix MMCs reinforced with graphite and boron fibers and SiC and Al2O3 particles with a lesser amount of data on magnesium and titanium MMCs. Hihara and Latanision (1) have summarized the corrosion behavior of metal matrix composites. However, there is relatively little data on the stress corrosion and corrosion fatigue performance of MMCs. 375 376 Jones II. CORROSION BEHAVIOR A. Discontinuous Reinforcement 1. Aluminum Matrix Composites The addition of a nonconducting reinforcing phases such as SiC or Al2O3 particles do not induce galvanic currents but can induce pitting corrosion. Trzaskoma et al. (2) and Trzaskoma (3) measured the pitting behavior of aluminum alloys 2024, 6061, and 5456 reinforced with 20 vol% SiC whiskers with 0.5–1.0-µm diameters and lengths up to 50 µm. Pitting potentials were determined electrochemically in 0.1N and 0.6N NaCl with and without dissolved oxygen. The pit initiation is unaffected for the 6061 and 5456 Al alloys by the presence of the SiC whiskers, and the 2024 Al alloy was affected by their presence. SiC whiskers caused a 100-mV shift in the pitting potential of the 2024 Al alloy in the de-aerated 0.1N NaCl solution. The pit morphology of the 6061 Al alloy was changed from large irregularly shaped pits in the unreinforced material to shallower round pits in the SiC reinforced material. Therefore, although SiC whiskers did not affect the pitting potential in the 6061 Al alloy, it did alter the pitting morphology. The relatively small 100-mV shift in the pitting potential of the 2024 Al alloy is not too significant because this alloy has a higher potential potential than the 6061 and 5456 Al alloys. Shimizu et al. (4) also found that pit initiation was unaffected by the presence of 10 vol% SiC whiskers, but that once a pit was initiated, the possibility for accelerated corrosion at the crevices between the reinforcement and the matrix existed. The electrochemical polarization curves for both 10 and 20 vol% SiC whisker reinforced 6061 Al in 3.5% NaCl at 25°C showed no shift in the open-circuit potential but an increased anodic current density. In contrast, SiC whisker reinforcement had no affect on the anodic polarization of 7075 alloy tested in the same solutions; however, in both alloys, the cathodic current density was greater for the reinforced material relative to the unreinforced. The authors suggested that this increased cathodic current could be the result of an interfacial layer between the matrix and SiC whiskers. Paciej and Agarwala (5) also examined the corrosion behavior of a SiCwhisker-reinforced (20 vol%) aluminum alloy. Their matrix alloy was AA 7090, which is an age-hardenable alloy with high Zn (6%) with Mg and Cu and is therefore a different matrix composition than the 2024, 5456, and 6061 Al alloys studied by Trzaskoma and co-workers (2,3). Paciej and Agarwala (5) also studied the effect of heat treatment on the corrosion behavior of their composite material that was a powder metallurgy alloy MA-87 of the same composition as the AA 7091 for comparison. A primary conclusion of their study was that the skin of the processed plate exhibited different corrosion behavior than the core in 3.5% NaCl. This difference was suggested as being due to Fe enrichment and increased porosity in the skin and Cu enrichment in the core. The steady-state corrosion Metal Matrix Composites 377 potential of the core region was about 100 mV anodic as compared to the MA87 material. Paciej and Agarwala (6) found that AA 7091 reinforced with particulate SiC (20 vol%) did not exhibit the same behavior as the whisker-reinforced material. Clusters of particles were observed to serve as regions for localized corrosion, but there was no difference between the skin and the core. The pitting potential in 3.5% NaCl ranged from 100 mV to no difference for the composite material relative to the MA-87, depending on the heat treatment. The small 100mV shift is similar to that observed by Trzaskoma and co-workers (2,3) for SiCwhisker-reinforced AA 2024. The effects of variable volume fraction (15–40%) of SiC particles on the corrosion behavior of AA 6061 was studied by Sun et al. (7). Studies were conducted in aerated and de-aerated solutions of NaCl. These authors found no shift in the corrosion potential with respect to the volume fraction of SiC, but they did see increased corrosion rates and less stable films with increasing volume fraction of SiC. A significant density of small pits was observed for NaCl concentrations ranging from 0.035% to 3.5% for potentials beyond the pitting potential, and these pits led to exfoliation-type cracking. Nunes and Ramanathan (8) compared the corrosion behavior of SiC (5 and 10 vol%) and Al2O3 (5 and 20 vol%) particle-reinforced aluminum matrix material. The matrix alloys were Al–7.5%, Si–1% Mg, and AA 2014 and the composite was produced by a melt-stirring process rather than a powder process as that reported by Trzaskoma and co-workers (2,3), Paciej and Agarwala (5,6), and Sun et al. (7). The matrix composition of the Nunes and Ramanathan (8) material is also different from the previous studies, with the exception of the AA 2024 reported by Trzaskoma and co-workers. Both types of reinforcement caused increased corrosion during immersion testing in 3.5% NaCl. The increased corrosion rate was associated with pits or microcrevices near the particle–matrix interface and from particle dropout. Pits in the SiC-reinforced material were deeper than in the Al2O3-reinforced composite. The pitting potential of the composite materials was shifted to lower potentials compared to the alloy, with this shift ranging from 95 mV in the aerated solution to 119 mV in the de-aerated solution. These shifts were similar for both particles, although there is some uncertainty in the data, and were similar in magnitude to that observed by Trzaskoma and co-workers (2,3) and Paciej and Agarwala (5,6). Further information on the effect of SiC particles on the interfacial corrosion of an Al 2024 alloy in a NaCl solution was presented by Yao and Zhu (9). They noted that the interfacial preferential dissolution (IPD) zone size was similar to the plastic accommodation size around the particles. They concluded from this that IPD was caused by the low integrity of the particle surface oxide film and not by chemical, metallurgical, or galvanic coupling effects. The corrosion characteristics of a cast Al–Si alloy containing 3 wt% graphite was evaluated by Saxena et al. (10). They measured the weight change follow- 378 Jones ing immersion in 3.5% NaCl and a salt spray and noted that the composite showed a factor of 16 for the immersion test and 6 for the salt spray test at 25°C. Potentiodynamic polarization tests in 3.5% NaCl showed considerably higher anodic and cathodic current densities for the composite relative to the base material. The authors concluded that the galvanic coupling of the graphite to the aluminum could explain the increased corrosion rates of the composite relative to the matrix material. They also observed localized corrosion at the interface between the graphite particles and the aluminum matrix as observed for composites with SiC and Al2O3 reinforcement. In summary, the primary affect of ceramic-phase reinforcement of aluminum alloys on corrosion is that of providing sites for pitting and crevice corrosion. This affect is manifested in decreased pitting potentials and increased corrosion rates from localized corrosion at particle–matrix interfaces and microcavities. Graphite particles have a much larger affect on the corrosion rate of aluminum composites than SiC or Al2O3 particle reinforcement because of the galvanic coupling and electrical conductivity of graphite. 2. Magnesium Matrix Composites Magnesium matrix composites offer attractive properties for aerospace and automotive applications because of their low density and excellent specific properties. Increases in the elastic modulus is one rationale for adding ceramic-phase reinforcements to magnesium. Magnesium has been reinforced with SiC (11,12) particles and Al2O3 (12) fibers, and because of the stability of SiC in magnesium, composites can be produced by conventional foundry processes. Luo (11) found good wetting between Mg and SiC, with no evidence of reaction at the Mg– SiC interface following liquid mixing and casting of the composite for the pure magnesium matrix composite. However, a reaction product of Mg2Si was present at the Mg–SiC interface in the AZ91 alloy matrix composite. The composite showed a 56% increase in the yield strength, but a decrease in toughness relative to the AZ91 alloy without reinforcement. Nunez-Lopez et al. (12) studied the corrosion behavior of a Mg–Zn–Cu (ZC71) alloy reinforced with 12 vol% SiC particles. The ZC71 alloy is not as corrosion resistant as the AZ91 series, so this composite would require a coating, as would the unreinforced matrix material. The composite was made by blending the SiC particles with the molten metal and then cast. Salt spray and electrochemical measurements were made in a 3.5% NaCl solution at 25°C. The salt spray corrosion rates for the composite material in the as-extruded and T6 condition were very similar to that of the unreinforced material in the same heat-treatment condition. However, localized corrosion is dominant in both reinforced and unreinforced material, and the local penetration rate for the reinforced material is about three times faster than the unreinforced material. This alloy develops a Metal Matrix Composites 379 passive film at cathodic potentials in both the reinforced and unreinforced condition and there is no evidence for galvanic corrosion between the matrix and SiC particles. The electrochemical polarization response of the reinforced and unreinforced material were nearly identical. Reinforcement of AZ91, the most corrosion-resistant magnesium alloy, with Al2O3 (13) fibers produced a shift in the electrochemical polarization response for high-pH (10.5) solutions. There was a shift of about 150 mV of the open-circuit potential in the anodic direction for the reinforced material relative to the unreinforced material with and without NaCl addition (3.5%). The shift in the open-circuit potential with Cl⫺ concentration was similar for both reinforced and unreinforced material, whereas the corrosion rates were similar up to about 1% NaCl. The corrosion rate of the reinforced material increased much more rapidly than the unreinforced material for Cl⫺ ion concentrations greater than 1%. The corrosion rate for the composite was three times greater than the matrix material at 3.5% NaCl. B. Continuous Reinforcement 1. Aluminum Matrix Composites Continuous fiber reinforcement offers a number of advantages over discontinuous reinforcement of metals. Most of these advantages are in the mechanical properties such as specific strength and stiffness. Specific strengths of 500 MPa/mg/ m3 and specific stiffness of 80 GPa/mg/m3 are possible with continuous-fiberreinforced aluminum (14) where only graphite/epoxy has higher values. Graphite, boron, and alumina fibers have been used to reinforce aluminum, but the transverse properties and corrosion performance of continuous-fiber composites are not as good as those of discontinuous reinforced composites. Some other continuous reinforced composites include unidirectionally solidified eutectics such as Al–Al4Ca (15) and metal–polymer laminates (16). Continuous-fiber-reinforced composites have been characterized by Trzaskoma (17) as Type I, where the reinforcing fibers are exposed on four of six sides of a cube, and discontinuous reinforcement as Type II where the reinforcement is exposed on all six sides of a cube. Crevice corrosion, galvanic corrosion, and pitting are possible in both Type I and II configurations; however, continuous fibers offer the potential for deep crevices in the direction of the fibers and graphite fibers produce significant galvanic effects. Both Pohlman (18) and Sedriks et al. (19) concluded that localized corrosion of B–Al composites was not the result of galvanic effects. Pohlman (18) evaluated B–Al couples and found no galvanic current, whereas boron fibers extracted from a composite that had a B–Al intermetallic on the surface did show galvanic effects. Therefore, a galvanic effect is expected for hot-pressed B–Al 380 Jones composites where this intermetallic formed at the fiber–matrix interface. Pohlman (17) associated the localized corrosion observed at the fiber–matrix interface with the presence of the B–Al intermetallic phase. Sedriks et al. (19) measured the corrosion behavior of B-fiber-reinforced 2024 aluminum alloy. They reported the loss in tensile strength of 2024 aluminum reinforced with 15% and 40% B following heat treatment for various times at 190°C and exposed with and without stress to a solution of 53 g of NaCl per liter of water. The maximum strength loss occurred for aging treatments of 1–2 h, with the matrix losing 45%, the 2024–15% B composite losing 50%, and the 2024–40% B composite losing 60%. Specimens exposed to this same environment but with a stress of 90% of the yield strength again showed an effect of the volume fraction of B reinforcement where specimens with the matrix material did not fail in 4000 h, 2024–15% B specimens failed in 1000 h, and 2024–40% B specimens failed in 40 h. Evans and Braddick (20) noted a preferential attack at the fiber–matrix interface for Bfiber-reinforced aluminum composites. They concluded that this localized corrosion was likely the result of an Al boride that formed at the interface during hot pressing. Weight-loss results show a rapid increase with time over the first 80 h, reaching a limiting weight loss of about 5 ⫻ 10⫺3 g/cm2. Similar results were obtained for a C-fiber-reinforced aluminum composite. One of the earliest reports of the corrosion behavior of C–Al composites was that by Dull et al. (21) for 6061 aluminum reinforced with carbon fibers. These materials were made by laying up alternating layers of fibers and foils and hot pressing to achieve the desired density. Corrosion tests were conducted in 3.5% NaCl at temperatures between 25°C and 75°C. The corrosion rate of the composite was 15 times greater than that of the alloy when tested at 25°C and increased substantially for temperatures exceeding 50°C. A graphite fiber–6061 Al composite was also evaluated for its corrosion performance by Aylor et al. (22), although the material was produced by infiltrating graphite fiber tows with 6061 aluminum alloy and hot pressing these tows. The graphite–6061 Al tows were hot pressed between 0.3-mm-thick alloy foils so that the outer skin had no exposed fibers, although the edges did have exposed graphite as in the Type I configuration discussed by Trzaskoma (16). Aylor et al. (22) also found that the composite exhibited considerably faster corrosion than the matrix material in splash and spray tests, as did Dull et al. (21). Pitting of the 6061 Al outer foil, which was the same on the composite and unreinforced 6061 Al, preceded the rapid corrosion of the underlying composite. Galvanic effects of the graphite– Al couple were suggested as the primary cause of the accelerated corrosion. Shimiziu et al. (4) noted that the pitting potential for graphite-fiber-reinforced 6061 Al was the same as the unreinforced alloy but that pit initiation was suggested as being easier because of the galvanic couple between graphite and the 6061 Al. They also noted a 100-mV shift of the open-circuit potential in the Metal Matrix Composites 381 anodic direct and a much higher cathodic current density than the unreinforced material or material reinforced with SiC whiskers or Al2O3 fibers. In studies to evaluate the corrosion mechanism in fiber-reinforced aluminum composites, Hihara and Latanision (23,24) found residual chloride in a graphite–6061 Al composite and measured galvanic corrosion for SiC, C, and TiB2 fibers coupled to 6061 Al. The chloride was traced to the aluminum infiltration process and was suggested as a factor in accelerated corrosion of these composites. The galvanic current density for the graphite fiber–aluminum couple was 30 times that of SiC or TiB2 fiber–aluminum couples. The galvanic current density was controlled by the rate of O2 reduction on the graphite and, therefore, the galvanic current density was greater in aerated solution than nonaerated solutions. In an effort to develop a corrosion-resistant graphite–aluminum composite, Wendt et al. (25) evaluated the corrosion performance of graphite-reinforced Al– Mo alloys. Alloys with 11–23% Mo were produced by sputter deposition and the corrosion behavior measured in 0.1M NaCl at a pH of 8 and a temperature of 25°C. Galvanic current densities for the Al–18% Mo and Al–23% Mo alloys coupled to graphite were three orders of magnitude less than for pure Al coupled to graphite. A directionally solidified eutectic of Al reinforced with lamina of Al4Ca intermetallic (15) can be considered in the same category as continuous-fiberreinforced metal matrix composites because of the potential for continuous crevice corrosion in the direction of the Al4Ca intermetallic. Corrosion testing was conducted in 3.5% NaCl at pH of 7 and a temperature of 25°C. The electrochemical response of the eutectic was essentially equal to that of the extruded alloy while being better than an Al–Cu alloy. 2. Magnesium Matrix Composites Magnesium–graphite composites with very high specific strengths and stiffnesses and low coefficients of thermal expansion are attractive materials for use in space applications. Although these materials could perform very well in the dry vacuum of space, they must exhibit sufficient corrosion resistance to survive manufacture, storage, and transport on Earth in the presence of moist salt air. Magnesium is one of the most reactive metals, so the galvanic effects between graphite and magnesium is of considerable importance. Trzaskoma (26) evaluated the corrosion performance of graphite-fiber-reinforced AZ91C which had AZ31B outer foils. Open-circuit potential-time and galvanic short-circuit current measurements were conducted in a borated boric acid solution containing 1000 ppm of NaCl at a pH of 8.4. Severe localized corrosion occurred at the edges where there was exposed graphite fibers after 5 days of immersion in this solution. This corrosion was evident (a) along the face plate, (b) in the underlying metal, and (c) as damage 382 Jones to the fibers. The open-circuit potential shifted with time in a very similar manner for the reinforced and unreinforced AZ91C material. A galvanic corrosion rate for the exposed edges of this composite was determined to be 1.8 mg/cm2 /day, which is very fast for structural materials. A potentially more stable Mg matrix composite reinforced with SiC fibers was evaluated by Hihara and Kondepudi (27). The composite was a unidirectionally aligned SiC monofilaments in a matrix of ZE41 aged at 329°C for 2 h to a T5 temper. The volume fraction of SiC monofilaments was about 50%. Corrosion studies were conducted in near-neutral 0.5M NaNO3 solutions at 30°C, de-aerated with high-purity N2. The magnesium alloy ZE41 had the following composition: 4.2% Zn, 0.7% Zr, 1.2% rare earth. Hihara and Kondepudi (27) found that the presence of the SiC monofilament caused the MMC to corrode faster in aerated than de-aerated solutions, unlike the matrix alloy in which the corrosion behavior was insensitive to oxygen concentration. This difference in corrosion behavior was evidenced by a 100-mV shift in the open-circuit potential of the composite material in aerated solutions relative to de-aerated solutions. The effect of O2 was explained by the reduction of O2 on the SiC. The corrosion rate of both the matrix material and the composite were very high at small voltages anodic to the open-circuit potential. A corrosion current density of about 3.2 ⫻ 10⫺3 A/ cm2 was reported for the composite. 3. Titanium Matrix Composite Titanium reinforced with monofilament SiC have potential application in drive shafts, turbine engine disks, compressor disks, and hollow fan blades and were being considered for the skin of the National Aerospace Plane prior to the cancellation of this program. The most common matrix materials are Ti–6Al–4V and Ti–15V–3Cr–3Sn–3Al (Ti-15-3) and the most common fiber is the Textron Specialty Materials SiC. Hihara and Tamirisa (28) examined the corrosion resistance of Ti-15-3 reinforced with monofilament SiC fibers. The composite had nine plies and was unidirectionally reinforced with SCS-6 fibers to volume fraction of 40%. The composite was given a heat treatment of 790°C for 30 min and aged at 510°C for 8 h. The Ti-15-3 matrix exhibited passive behavior in a solution of 3.15 wt% NaCl, whereas the composite material did not. The current density was at least a factor of 10 greater for the composite. Hihara and Tamirisa (28) calculated the polarization behavior of the composite from corrosion measurements conducted separately on the matrix and fibers and calculated using mixed potential theory. The experimental and calculated polarization curves were identical, which led the authors to conclude that the composite fabrication process had not altered the electrochemical behavior of the constituents. At open-circuit potential, the composite material exhibited excellent corrosion resistance. Metal Matrix Composites 383 III. STRESS CORROSION CRACKING AND CORROSION FATIGUE BEHAVIOR Metal matrix composites have the potential to exhibit stress-corrosion properties superior to those of the base alloy because the reinforcement provided by particles, whiskers, or fibers reduces the crack tip strain rate. Of course, this potential can only be realized if the corrosion performance of the composite is not degraded relative to the matrix material. As discussed earlier, this is most likely the case for chemically stable, nonconducting, ceramic reinforcements such as SiC or Al2O3. The localized corrosion that has been noted for many of the MMCs could impact the stress-corrosion or corrosion fatigue performance of MMCs. The following sections summarize the crack growth observations for MMCs with discontinuous and continuous reinforcements. Much of the data are for cyclically loaded specimens because stress-corrosion measurements are more difficult because of the high K Iscc /K Ic ratio. A. Discontinuous Reinforcement Some of the first measurements of the effect of particle and whisker reinforcement on corrosion fatigue were made by Yau (29) and Hasson et al. (30). Yau measured the fatigue crack growth behavior of the 6061 Al–SiC composite in air, water, and water with 3.5% NaCl. The behavior in water with and without NaCl was identical but exhibited slightly greater crack velocities than the tests in air. This suggests that water and water plus NaCl likely caused some increased corrosion rate and degradation of the composite relative to the matrix material. However, the reinforced material exhibited a slower crack velocity in water plus 3.5% NaCl than the unreinforced material. Therefore, if the localized corrosion effect of the reinforcement can be eliminated, these results suggest the possibility for a beneficial effect of reinforcement on stress corrosion and corrosion fatigue. Hasson et al. (30) evaluated unreinforced 6061-T6 and 6061-T6 reinforced with 20% SiC whiskers and particles. The l/d ratio for the whiskers ranged from 2 to 5 following thermal–mechanical processing. Their results show the cycles to failure as a function of alternating stress in laboratory air and moist salt air. The whisker-reinforced material had 100 times greater cycles to failure than the unreinforced material when tested in moist salt air at an alternating stress of 96 MPa. This is a significant improvement imparted by the presence of the SiC whiskers. Hasson et al. (30) did not report corrosion results for their test conditions, but it is likely that moist salt air caused less localized corrosion and degradation than aqueous corrosion tests where localized corrosion has been noted for MMCs. The improvement in corrosion fatigue imparted by the SiC whiskers was reduced at high alternating stresses. For instance, at an alternating stress of 144 MPa, the whisker reinforcement imparted less than a factor of 10 increase in 384 Jones cycles to failure. Also, in contrast to the observations of Yau (29), the cycles to failure of the reinforced material tested in moist salt air exceeded the cycles to failure of the unreinforced material tested in laboratory air. Corrosion fatigue results of 6061 Al reinforced with 10% Al2O3 and tested in air and 3.5% NaCl have recently been reported by Bertolini et al. (31). They observed that there was little difference in the da/dn versus ∆K curve between air and 3.5% NaCl for material extruded and forged 17%, but there was an increase in da/dn for material extruded 8.3%. However, the increase in da/dn was only about a factor of 5. Pitting at particle–matrix interfaces was noted following potentiodynamic corrosion tests, as noted by others. The authors concluded that the pitting corrosion contributed to the corrosion fatigue crack growth results; however, they did not evaluate an unreinforced 6061 Al alloy to determine if the particleinduced localized pitting was a significant factor in the da/dn versus ∆K behavior of the composite. B. Continuous Reinforcement 1. Aluminum Matrix Composites Experimental results on the stress corrosion or corrosion fatigue of continuousor fiber-reinforced material is equally sparse as it is for discontinuous particle or whisker-reinforced material. Berkeley et al. (32) measured the remaining strength of 6061 Al reinforced with Nextel 440 fibers after exposure to NaCl at pH levels of 1.5 and 2.0. Nextel 440 has a composition of 70% Al2O3, 28% SiO2, and 2% B2O3. The composite contained 45% Nextel fibers, and the specimens were exposed to the ASTM G44 solution stressed to 80% of their yield strength. The strength of the composite following exposure was used as a measure of the environmental degradation of the composite. Although this method may measure the effects of stress-corrosion cracking by cracks that grow into the material, it may only be measuring an embrittling effect. Preferential corrosion at the fiber–matrix interface following the length of the fiber was observed following exposure to this low-pH solution and was connected to the strength reduction observed. The residual strength of the composite material following 5 days of exposure to the ASTM G44 solution was about 80% regardless of whether a stress was applied during the exposure. For the same exposure conditions, the 6061 Al alloy had a residual strength of 95%. The authors concluded that localized corrosion along the length of the fiber at the fiber–matrix interface was the primary cause of the larger strength reduction for the composite compared to the matrix alloy. Sedriks et al. (33) evaluated diffusion-bonded 2024 aluminum alloy–boron filament composites and reported localized corrosion of the transverse ends with and without an applied stress. Stress-corrosion tests of samples loaded in the transverse direction in a NaCl solution to 90% of the yield strength failed in 1000 Metal Matrix Composites 385 h for 15% B fibers and 40 h for 40% B fibers. The failures were intergranular in the matrix and interfacial between the matrix and fibers. No failures were observed for times up to 4000 h in the 2024 alloy tested in the longitudinal direction. Stress-corrosion tests of samples loaded in the longitudinal direction showed that the composite outperformed the 2024 Al in NaCl solution. No failures occurred in the composite for times up to 1000 h, whereas the matrix failed in less than 10 h. The inferior stress-corrosion properties of the composite in the transverse direction resulted from the exposed transverse cross sections. A unique composite comprised of metal matrix and polymer composite laminates has been evaluated by Wanhill (16). Corrosion fatigue tests were conducted on 2024 Al-T3 with carbon–epoxy laminates and 7475 Al-T761 with carbon–epoxy laminates in air and air with a water spray. The specimens were a symmetrical laminate of Al face sheets over a unidirectional carbon fiber–epoxy composite core. These materials were being evaluated for aerospace applications so they were loaded with a flight simulation loading program with a mean load of about 20% of the ultimate load capacity of the composite. The possibility of a galvanic potential between the carbon fibers in the epoxy composite and the aluminum alloys existed, but there was no evidence for accelerated corrosion from galvanic corrosion. The results are reported in number of flights to produce a crack extension of 40 mm. In the air plus water spray tests, the 2024-T3 carbon– epoxy composite was superior to the 2024-T3 material and the other composites. In an environment of air plus water spray, the 2024-T3 carbon–epoxy composite required 8300 flights for 40 mm of crack extension, whereas the matrix required only 4800 flights and the 7475-T761 carbon–epoxy composite only 1800 flights. 2. Magnesium Matrix Composite Stress-corrosion tests of a magnesium alloy reinforced with alumina fibers reported by Evans (34) is possibly the strongest case for the potential benefit of composite reinforcement on stress corrosion and corrosion fatigue. A magnesium alloy, ZE41A, with a composition of 4.5 Zn, 0.7 Zr, and 1.0% rare earths was reinforced with 55 vol% alumina fibers. Tests were conducted with fibers oriented perpendicular (90°) and parallel (0°) to the principal stress in a solution of NaCl and K2CrO4 and reported as time to failure versus applied stress and crack velocities (determined from the change in the bending moment with time) versus time. Both unnotched and chevron-notched specimens were loaded in a cantilever configuration. For the unnotched specimens, the failure time was measured as a function of the applied stress and a clear benefit of the alumina fibers was noted for the 0° composite. Composites with the 90° fiber orientation had properties similar to that of the matrix. A threshold stress of 600 MPa was found for the 0° composite and about 200 MPa for both the matrix and 90° composites. For the chevronnotched specimens, a crack growth velocity of about 10⫺5 mm/s was estimated for 386 Jones the 0° composite. However, the K th for the 90° composite was only 6.5 MPa⋅m1/2 while it was 16.2 and 11.2 MPa⋅m1/2 for the 0° composite and the matrix, respectively. 3. Titanium Matrix Composites Mahulikar and Marcus (35) evaluated the fatigue crack growth behavior of B4C– B and BORSIC reinforced Ti–6Al–4V in humid air. Fatigue tests were conducted at an R of 0.1 to measure crack closure effects and compared to results at an R of 0.5 and with variable water-vapor pressures. The primary observation was that a transition from interface separation to fiber splitting occurred at a water-vapor pressure of 100 MPa. The fiber splitting was explained as resulting from crack closure effects rather than a direct embrittlement of the fiber by the environment. Increased crack growth velocities accompanied the occurrence of fiber splitting. Ti–6Al–4V/carbon–epoxy laminate composites were also evaluated by Wanhill (16) along with the 2024 Al and 7475 Al laminate composites described earlier. The Ti–6Al–4V/carbon–epoxy composite required only 1400 flights for a crack extension of 40 mm when tested in an air plus water spray environment and about 3600 cycles in air. This is in contrast to 8300 flights for the 2024 Al/ carbon–epoxy and 4800 flights for the 7475/carbon–epoxy composites. The poor performance of the Ti–6Al–4V was related to cracking of the carbon–epoxy layer. 4. Iron Matrix Composite The stress-corrosion performance of a laminate composite of maraging steel and Armco iron was determined by Floreen et al. (36). The laminate was produced by hot-rolled bonding plates of the 18% Ni maraging steel with a sheet of Armco iron between each plate. Tests were performed with a notch-bend specimen with the notch oriented in a notch-arrester orientation (cracks running through the thickness of the plates) and crack-divider orientation (cracks running parallel to the layer interfaces). Composites with two, four, and eight layers of maraging steel were tested. Tests were performed in an aerated 3.5% NaCl solution. In all cases, the composite failure time was over 240 h for the crack-arrester orientation but ranged between 63 and 80 h for the crack-divider orientation for two, four, and eight layers of maraging steel. Stress-corrosion cracks in the crack-arrester specimens propagated through the maraging steel but were diverted and branched in the Armco iron. The Armco iron layer was an effective crack-arrest layer. 5. Model of Stress-Corrosion Cracking in Composite Materials A model of the stress-corrosion crack growth rate of metal matrix composites has been presented by Jones (37). This model considers the length-to-diameter Metal Matrix Composites 387 (l/d) ratio and volume fraction of reinforcing phase and the matrix stress exponent. The model is based on the assumption that several stress-corrosion mechanisms are controlled by crack tip strain rate and that the presence of a reinforcing phase reduces the crack tip strain rate. Processes such as the passive film rupture rate, dislocation–surface interaction rate, surface reactivity rate, and crack opening effects on transport of species into and out of the crack and crack wall corrosion rates are all factors that depend on the crack tip strain rate. The resulting predictions of this model are presented in Figs. 1 and 2, along with data by Yau (29), Hasson et al. (30), and Jones (37). The model predictions given in Fig. 1 show that the value of R, the ratio between the stress corrosion velocity of the composite to that of the unreinforced matrix, is a strong function of the l/d ratio and moderately dependent on n, the matrix creep stress exponent. Increasing values of the matrix stress exponent causes a decrease in the value of R. The predictions given in Fig. 2 show that R has only a small dependence on the volume fraction of the reinforcing phase over the range 5–40%. This model assumes that the reinforcement does not accelerate the corrosion rate and induce other damage. As discussed earlier, discontinuous reinforcement with ceramic phases often cause localized corrosion around clusters of particles or the particle– matrix interface. This localized corrosion very likely influences the quality of the Fig. 1 Calculated and experimental values of R versus l/d and n for 7090 Al/SiC and 6061 Al/SiC. 388 Fig. 2 Jones Calculated values of R versus l/d and Vf. comparison between the model and experiment. The results of Jones (37) and Yau (29) are consistent with the model and matrix stress exponent of 3–6. Hasson et al. (30) results are plotted at an l/d of 100, which is based on the whisker dimensions prior to processing the composite, but they reported that the l/d was reduced to about 5 after processing. Therefore, their results should be shifted substantially to the left as the arrow indicates. IV. CORROSION PROTECTION Many of the desirable properties of metal matrix composites are compromised by corrosion, especially for continuous reinforcement with graphite. Because of this, several schemes have been evaluated to provide corrosion protection of these materials. Aylor et al. (38) evaluated the effects of electroless nickel coatings and anodization on the corrosion performance of 6061 Al with continuous graphite fibers. Tests were performed by immersion in seawater, splash, spray, and atmospheric exposure. Corrosion on the uncoated composite began with pitting and progressed to galvanic corrosion between the graphite fibers and the aluminum matrix. Electroless nickel coating resulted in severe galvanic corrosion of the 6061 Al–graphite substrate because of flaws in the coating and the high Ni/ Al ratio, and therefore cathode-to-anode ratio, at these flaws. Anodization increased the corrosion resistance of the composite with some thinning of the anod- Metal Matrix Composites 389 ized layer that did not impact the corrosion resistance for times up to 180 days. Anodization was judged to be a useful method for protecting metal matrix composites. Mansfeld et al. (39) concluded that chromate conversion coatings and immersion in CeCl3 both produced stable, protective coatings on a SiC–Al composite. Cathodic protection is another route for protecting these composites, but Hihara and Latanision (40) have demonstrated that cathodic overprotection can lead to accelerated corrosion for aluminum matrix composites. This occurs because of the production of OH⫺ and the instability of aluminum in high-pH environments. The effect of cathodic overprotection was the most severe in the 6061 Al–graphite composite and less so in the 6061 Al/SiC composite and 6061 Al. Modifying the matrix corrosion behavior through alloy compositional changes is a route to improving the corrosion performance of Al–graphite composites considered by Wendt et al. (25). Aluminum and Al–Mg alloys with 11%, 18%, and 23% molybdenum were prepared by sputter deposition and the galvanic corrosion response of these alloys with graphite determined from galvanic diagrams constructed from the polarization data. The authors concluded that an Al–18% Mo/graphite composite would exhibit good galvanic corrosion behavior. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. LH Hihara, RM Latanision. Int Mater Rev 39(6):245, 1994. PP Trzaskoma, E McCafferty, CR Crowe. J Electrochem Soc 130:804, 1983. PP Trzaskoma. Corrosion 46:402, 1990. Y Shimizu, T Nishimura, I Matsushima. Mater Sci Eng A 198:113, 1995. RC Paciej, VS Agarwala. Corrosion 42:718, 1986. RC Paciej, VS Agarwala. Corrosion/87. Houston, TX: National Association of Corrosion Engineers, 1987, Paper 221. H Sun, EY Koo, HG Wheat. Corrosion 47:741, 1991. PCR Nunes, LV Ramanathan. Corrosion 51:610, 1995. H-Y Yao, RZ Zhu. Corrosion 54:499, 1998. M Saxena, OP Modi, AH Yegneswaran, PK Rohatgi. 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INTRODUCTION Ceramic matrix composites (CMCs) are being developed to take advantage of the high-temperature properties of ceramics while overcoming the low fracture toughness of monolithic ceramics. Toughening mechanisms, such as matrix cracking, crack deflection, interface debonding, crack-wake bridging, and fiber pullout, are being incorporated in CMCs to reduce the tendency for catastrophic failure found in monolithic ceramics. Ceramics reinforced with particulate, whiskers, and continuous fibers exhibit varying aspects of these toughening mechanisms; however, reinforcement with continuous fibers offers the greatest improvements in toughness. Composites with carbide, oxide, glass, and carbon matrices are being utilized in the development of CMCs. In the case of carbide, oxide, and glass matrix CMCs, the matrix exhibits excellent high-temperature corrosion resistance so that a goal of the composite development is to not detract from this preexisting property. This is not the case for carbon matrix composites, which frequently need coatings to provide adequate corrosion protection. The purpose of this chapter is to review the database and understanding of corrosion behavior of CMCs with the intent that this information will be useful in the development of materials with improved performance and reliability. Composite materials are chemically and microstructurally heterogeneous, consisting of matrix, matrix–reinforcement interface, and reinforcement constituents. The corrosion behavior of each of these constituents will likely not be equal whether evaluated individually or within the composite. The composite corrosion resistance may be more complex than each constituent because of inter391 392 Jones et al. actions between corrosion reaction products and the composite constituents. An example of this interaction would be a reaction product from the corrosion of the matrix that protects a less corrosion-resistant reinforcement or reinforcement– matrix interface. However, there is far more data on the corrosion behavior of the constituents as monolithic materials than there are on the composites made from these constituents. It is important to note that the constituents of a composite may differ slightly or substantially in chemical composition, crystal structure, and microstructure from their monolithic equivalent. Examples of these differences are the SiC fibers and matrix produced from polycarbosilane or its variants that may contain considerable oxygen or may be nanocrystals embedded in an amorphous matrix. There is considerable value in evaluating the corrosion resistance of the monolithic equivalent of the composite constituents because these generally represent the baseline corrosion behavior. An effort will be made to identify where significant differences are likely to occur between the composite constituents and the monolithic equivalent. High-temperature composites are being developed to operate in a variety of environments containing alkali elements, mildly oxidizing mixed gases such as He⫹O2, highly oxidizing environments, and H2. Not all ceramic composites are being considered for each of these environments, so the corrosion data are not available for each ceramic composite in every environment. II. SILICON CARBIDE MATRIX COMPOSITES A. Corrosion Reactions for Silicon Carbide Silicon carbide will chemically react with O2, H2, and H2O according to the following reactions: SiC(s) ⫹ O2(g) ⫽ SiO(g) ⫹ CO(g): low pO2 (1a) SiC(s) ⫹ / O2(g) ⫽ SiO2(s) ⫹ CO(g): high pO2 (1b) 32 SiC(s) ⫹ 2H2(g) ⫽ Si ⫹ CH4(g) (2) SiC(s) ⫹ 2H2O(g) ⫽ SiO(g) ⫹ CO(g) ⫹ 2H2(g) (3) Silicon carbide is thermodynamically unstable in O2, H2, and H2O environments under certain conditions. However, the kinetics of these reactions can be affected by the formation of a protective layer of SiO2 at high pO2 such that SiC is stable in many corrosive environments. The stability of the passive layer then becomes critical to the stability of SiC. The O2 pressure for the SiC active–passive transition was determined by Gulbransen and Jansson (1) as shown in Fig. 1. At temperatures of 800–1000°C, this transition occurs at O2 pressures of 10⫺8 atm. Therefore, the pertinent reactions then become those between SiO2 and specific gaseous Ceramic Matrix Composites 393 Fig. 1 Transition pressures for SiC active–passive oxidation versus temperature, according to Gulbransen and Jansson (1). or molten-salt environments, except in molten Li with a low O2 activity, where SiO2 is unstable. Some relevant reactions with SiO2 are below: SiO2(s) ⫹ H2(g) ⫽ SiO(g) ⫹ H2O(g) (4) xSiO2(s) ⫹ Na2SO4(P) ⫽ Na2OX x(SiO2)(P) ⫹ SO3(g) (5) where M is an alkali element such as Na and Li. A phase diagram for the Na2O– SiO2 system is shown in Fig. 2. Alkali elements such as Na and Li cause a breakdown of the passive SiO2 film by the formation of low-melting alkali silicates such as those which occur at 800°C in the Na2 –SiO2 system and 1024°C in the Li2O–SiO2 system. The eutectic temperature in the Li2O–SiO2 system is about 250°C higher than that in the Na2O–SiO2 system and, therefore, Li is expected to have less effect on the passive film on SiC at temperatures below 1000°C than Na. A summary of the behavior of SiC in gas–molten-salt environments as presented by McKee and Chatterji (2) is shown in Fig. 3. Passivation occurs at high pO2, and active oxidation (formation of gaseous SiO) occurs at low pO2. A basic salt or salt melt with a low pO2 at the salt–SiC interface will cause active corrosion, as depicted by reaction scheme (4) and (5). McKee and Chatterji suggest that SiC will not react with H2; however, more recent analysis by Herbell et al. (3) indicates that the equilibrium partial pressure of CH4 [Eq. (2)] is 10⫺4 394 Fig. 2 Jones et al. Phase diagram for the system Na2O–SiO2. atm at temperatures of 850–1400°C for a H2 pressure of 1 atm. While McKee and Chatterji measured the sample weight change in the test environment as a function of time in 1 atm of H2 at 900°C, they observed no reaction between SiC and H2 (scheme 1, Fig. 3). However, their gas may have had sufficient O2 or H2O to cause passivation. Herbell et al. (3) calculated the SiO(g) partial pressure for Eq. (4) with 1 atm of H2 containing 1 ppm of H2O to be 10⫺7 atm. Therefore, it would appear that a very small amount of H2O mixed with O2 to promote SiO2 formation would be sufficient to cause a significant reduction in the reaction rate of SiC. Ceramic Matrix Composites 395 Fig. 3 Possible modes of behavior of SiC in gas–molten-salt environments. Jacobson (4) evaluated the kinetics and mechanisms of the corrosion of sintered α-SiC in molten salts at 1000°C. In the reaction of Na2SO4 /O2 with SiC, the reaction occurred primarily in the first few hours with the formation of a protective SiO2 layer. This observation was demonstrated with results showing the total weight of the corrosion products/unit area reaching 6 mg/cm2 (SiO2 ⫹ Na2O–x(SiO2) after a few hours and remaining constant up to 20 h. Jacobson and Smialek (5) also noted that SiC is subject to pitting corrosion in molten salts. Pits occurred at structural discontinuities and bubbles formed during the formation of SiO2. Pitting corrosion is detrimental because it demonstrated that the passive SiO2 layer has been degraded and because the pits act as flaws resulting in reduced fracture strength. Corrosion measurements of SiC immersed in a crucible of molten Na2SO4 were conducted by Tressler et al. (6). Weight change results for both SiC and Si3N4 in 100% Na2SO4 at 1000°C are given in Fig. 4, where it is evident that considerable corrosion occurs for both materials, with SiC showing a higher corrosion rate than Si3N4. Both materials exhibited a weight loss because the reaction 396 Jones et al. Fig. 4 Weight change versus exposure time to Na2SO4 at 1000°C for several ceramic composites. (From Ref. 6.) product was removed prior to measuring the weight change. An in situ measurement would probably exhibit a weight gain up to the time at which reaction product spallation occurs. Corrosion data obtained on SiC–Si3N4 and SiC–SiC composite materials by Henager and Jones (7) is also presented in Fig. 4. The corrosion rate of the composite material was found to exceed that of the monolithic material for a Nicalon–SiC-reinforced, hot-pressed Si3N4 and a Nicalon– SiC reinforced CVI SiC. A SiC-whisker-reinforced, hot-pressed Si3N4 exhibited a similar corrosion rate to the unreinforced hot-pressed Si3N4. The activation energy for the corrosion of the fiber-reinforced material was 880 kJ/mol, whereas the whisker-reinforced matrix material had an activation energy of 1280 kJ/mol. The difference in corrosion rate and activation energy for the fiber-reinforced material was associated with preferential corrosion of the Nicalon fiber. The fiberreinforced material had a carbon interfacial layer between the fiber and matrix, and the whisker-reinforced material had a thin amorphous glass layer. However, the difference in corrosion behavior is not thought to be associated with the interfacial layer but with the structure of the Nicalon fiber following hot pressing. As produced, Nicalon has an amorphous/microstryalline structure while the hotpressed structure was crystallized with evidence for a Mg silicate in the grain boundaries. The source of the Mg was the MgO sintering aid used for the hotpressing process. Corrosion of the Nicalon fiber along the glass-enriched grain Ceramic Matrix Composites 397 boundaries probably accounted for the high corrosion rate of the Nicalon-fiberreinforced composite material. The corrosion rate of a SiC–SiC composite produced by the CVI process was also greater than that of monolithic SiC (Fig. 4). Microstructures of CVI SiC–SiC composites made by the ORNL process and the Du Pont process both exhibit porosity, which exists in materials made by this process. Composites containing fibers coated with C, BN, C–B4C–BN, and no coating were evaluated. The corrosion rate was highest for the material made with fibers with no coating and the BN coating, whereas composites made with C and C–B4C–BN coatings exhibited smaller corrosion rates. However, the corrosion rate of these materials at 900°C was considerably greater than that of monolithic material at 1000°C. Matrix corrosion contributed the majority of the weight change, and while some fiber corrosion was noted, although this was a relatively small factor compared to the matrix corrosion. The open porosity and penetration of the molten salt into the composite are considered the primary factors causing the high corrosion rates for the SiC–SiC composite. Fox et al. found similar results in burner rig corrosion studies of CMCs to those obtained in molten salt (8). The burner rig tests utilize a hot-combustion gas and would be expected to differ from a molten Li or He gas environments; however, molten deposits are one mechanism for corrosion in hot-combustion gas studies. Fox et al. evaluated Si3N4 reinforced with 30 vol% SiC whiskers and fibers. The method of procesing for the whisker-reinforced materials was not reported; however, it was probably hot pressed, whereas the fiber-reinforced material was reaction bonded and had a residual porosity of 30%. Tests were conducted for 40 h at 1000°C in an environment containing 2 ppm Na added as NaCl. Monolithic Si3N4 exhibited a 10% decrease in the fracture stress and the whisker-reinforced material a 35% decrease. The fiber-reinforced material exhibited an excessive corrosion rate, but no fracture studies were conducted. The high residual porosity in the reaction-bonded material and the free C from the carboncoated fibers were postulated as the cause for the high corrosion rate in the fiberreinforced material. In the molten-salt corrosion studies (8), the more porous CVI composite material also exhibited a higher corrosion rate than did the hot-pressed Si3N4 /fiber- or whisker-reinforced-material. Effects of corrosion on the mechanical integrity of SiC–SiC structural components is an important factor in the durability of components constructed of SiC/SiC. Decreases in the fracture strength would be possible by corrosion penetration along grain boundaries. Smialek and Jacobson (9) also found that moltensalt-induced pits were responsible for up to 50% of the observed strength reduction of SiC after 48 h at 1000°C in Na2SO4 –SO3. A relationship of fracture stress versus (pit depth)⫺2 illustrated the flaw-induced fracture relationship. Henshall et al. (10) demonstrated that combustion gases containing alkali salts also contribute to accelerated subcritical crack growth of monolithic SiC (Fig. 5). The crack 398 Jones et al. Fig. 5 Crack velocity versus stress intensity for monolithic SiC exposed to air and hotgas environments. (From Ref. 10.) velocity in the hot-combustion gas was several orders of magnitude faster at a 50 K lower temperature than in air. The primary mechanism for this accelerated crack velocity is penetration of alkali ions into the grain-boundary glass phase and the reduction in the viscosity of this phase. The reduced viscosity causes increased crack tip creep rates and creep damage. B. Stability of SiC in Mildly Oxidizing Environments Spear et al. (11) measured the oxidation rate of hexagonal α-SiC platelets with both (0001) C and (0001) Si faces. The only solid corrosion product observed was SiO2 [Eq. (1b)], with CO(g) formed at the SiC–SiO2 interface. The O2 pressure was 10⫺3 to 1 atm at 1200–1500°C, which is above the active–passive transition; thus, the presence of SiO(g) would not be expected. An activation energy of 120 kJ/mol was found for SiC, which is very comparable to the value of 112 Ceramic Matrix Composites 399 kJ/mol found for Si. Silicon nitride has a much higher activation energy of 486 kJ/mol because of the formation of a silicon oxynitride phase between the Si3N4 and the SiO2 phases. A growth rate of 1.1 nm/min was determined for SiC at 1300°C in 1 atm of O2. Oxygen diffusion was postulated as the rate limiting process up to 1350°C, above which ionic oxygen diffusion was rate controlling. The protective properties of the passive SiO2 film are very sensitive to the impurities of the SiC from which the film is formed. Comparison between the oxidation behavior of several batches of SiC to that of very high-purity SiC produced by chemical vapor deposition was made by Fergus and Worrell (12). They observed an activation energy of 142 kJ/mol for CVD SiC and 217 kJ/mol for sintered α-SiC. This compares with the value of 120 kJ/mol reported by Spear et al. for single-crystal SiC platelets. Both the CVD and SiC platelets are of much higher purity than the sintered α-SiC to which sintering aids were added. Vaughn and Maahs (13) in a review of the active–passive transition for SiC showed that this transition was quite variable depending on the source of the SiC and other experimental conditions such as the gas flow velocity. The Gulbransen data (Fig. 1) was very close to that determined theoretically, but most other materials exhibited much higher transition temperatures, presumably due in part to variable impurity concentrations in the materials. Results of oxidation studies of ceramic composites have not been reported, although Luthra (14) has presented some theoretical concepts on the oxidation of ceramic composites. Luthra assessed the gas-phase diffusion of oxygen through microcracks, solid-state diffusion of oxygen through protective oxide layers, formation of gaseous reaction products, and the combined reaction of the matrix reinforcement and oxygen. In particular, the diffusion of oxygen through microcracks to react with the reinforcing fiber or the interface between the fiber and matrix is a definite possibility for SiC–SiC composites produced by the CVI process because they have at least 10% residual porosity. Thermal cycling is also likely to induce matrix microcracking because of the thermal expansion mismatch between the reinforcing fiber and the matrix. For SiC fibers in a CVI SiC matrix, this mismatch will be small, but it is not zero. Luthra’s analysis is primarily for oxide matrix composites, where the oxygen will not react with the matrix, whereas for SiC where the oxygen could react with the matrix, the transport would be considerably slower. The structural behavior of CMCs in an inert gas environment will be dependent on (a) whether the material is undergoing active or passive oxidation, (b) the stability of the fiber–matrix interfacial layer in O2, and (c) the stability of the fiber in O2. Kim and Moorhead (15) have recently evaluated the flexural strength of α-SiC at room temperature following a 10-h exposure to Ar–O2 at 1400°C. Flexural strengths following exposure to Ar–O2 mixtures with O2 partial pressures of 7 ⫻ 10⫺7 to 2 ⫻ 10⫺4 MPa showed a strength reduction with increasing pO2 up to about 2 ⫻ 10⫺5 MPa and a complete restoration of the as-polished 400 Jones et al. strength at pO2 ⬎ 2 ⫻ 10⫺5 MPa (Fig. 6). The minimum strength (60% of unexposed material) at 2 ⫻ 10⫺5 MPa was the result of active corrosion creating strength-reducing flaws along grain boundaries and at pits. Easler et al. (16) found that the room-temperature flexure strength both increased and decreased upon exposure to air at 1370°C with and without an applied load. The strength increases were thought to result from oxidative blunting, whereas the decreases which occurred under load were due to subcritical crack growth producing larger flaws. Minford et al. (17) also found that stress enhanced the uptake of O2 into SiC. Oxygen penetration occurred along cation-enriched grain boundaries. A stress intensity threshold of about 1 MPa ⋅ m1/2, below which crack blunting occurred and above which crack extension occurred was found. Oxidation can alter the dynamic crack growth behavior of SiC as well as reduce the fracture strength due to corrosion-induced flaws. A reduction in the K th of α-SiC after being loaded to different stress intensities for 4 h at 1200°C and 1400°C was reported by Minford et al. (18). They reported a reduction in the K th from 2.25 in the unoxidized to 1.75 MPa ⋅ m1/2 in the oxidized conditions at 1200°C and from 1.75 to 1.25 MPa ⋅ m1/2 at 1400°C for samples tested in air or preoxidized. They concluded that the O2 caused a change from diffusioncontrolled crack growth to viscous cavity growth and linkage with a corresponding reduction in K th. The K th was defined as the transition between crack blunting Fig. 6 Room-temperature flexural strength of sintered α-SiC after exposure for 10 h at 1400°C in Ar atmospheres with various pO2. Ceramic Matrix Composites 401 and growth as determined from the fracture strength following the 4-h static loading at various K values. McHenry and Tressler (19) found that the K th and subcritical crack growth rate was independent of the pO2 for pressures ranging from 10⫺8 to 10⫺4 atm and temperatures of 900–1100°C. The crack growth rate exhibited an Arrhenius temperature dependence with an activation energy of 20 kcal/mol, which is consistent with viscous flow of a grain-boundary glass phase. The lack of a dependence on pO2 does not necessarily indicate that O2 did not induce crack growth but that the effect of O2 was saturated at pressures above 10⫺8 atm. C. Oxidizing Environments A key feature of the oxidation behavior of SiC–SiC in O2 at pressures of 2 ⫻ 104 Pa (atmospheric pressure of O2) observed by Windisch et al. (20) is that only a weight loss was observed. No SiO2 formation occurred in any of the materials with graphite interphases, although some boron-containing glass phase was observed for the material with a BN interphase. The kinetics of mass loss is shown as a function of pressure and temperature in Figs. 7 and 8. Complete burnout of the graphite interphase occurred in less than 104 s in the small test samples at a Fig. 7 Mass loss versus exposure time for SiC–SiC with a C interphase exposed to various O2 partial pressures at 1100°C. 402 Jones et al. Fig. 8 Mass loss versus time for SiC–SiC with a C interphase exposed at 2.5 ⫻ 103 Pa O2 at various temperatures. pressure of 2.4 ⫻ 104 Pa and a temperature of 1373 K. It is also clear that the reaction rate increases with increasing temperature. An activation energy of about 50 kJ/mol was reported by Windisch et al. (20), which could be explained as diffusion controlled through a boundary layer or as being reaction rate controlled. An interphase recession rate was determined from the weight-loss measurements and by direct physical measurement. Both methods gave very similar recession rate equations with the physically measured equation as follows: log(RR) ⫽ 0.9 log(pO2) ⫺ 9.9 (6) It should also be noted that only weight loss was observed over the temperature range 1073–1373 K, which borders on the temperature range (873–1073 K) suggested by Evans et al. (21) for the pest phenomena. Oxidation of SiC–SiC composites with carbon interphases can also result in the formation of SiO2 and a weight gain following an initial weight loss (22,23) or a reduced weight loss with increasing temperature (24). Tortorelli et al. (23) observed an initial weight loss followed by a weight gain in a Nicalon-fiberreinforced SiC–SiC composite with a 0.3-µm-thick graphite layer exposed to dry air (pO2 of 2 ⫻ 104 Pa) at 1223 K. Following the initial weight loss from oxidation Ceramic Matrix Composites 403 of the carbon, they found that SiO2 formation occurred within the interfacial region previously occupied by the carbon. For the Nicalon-reinforced composite material, complete carbon depletion occurred within 15 min at 1223 K, followed by weight gain from SiO2 formation. Unal et al. (24) observed their largest weight loss (5%) at 1223 K for an exposure of 50 h in dry O2 and a decreasing weight loss with increasing temperature up to 1673 K (2%). Their material was a Nicalonreinforced SiC–SiC with a 0.5-µm-thick fiber–matrix carbon interphase. Kleykamp et al. (25) observed the following reactions of air with SiC-fiber-reinforced SiC composites: (a) oxidation of free carbon at temperatures of 800–965 K and (b) a fast exothermic reaction and weight gain beginning at 1073 K and containing up to 1773 K and up to times of 1 h followed by (c) the diffusion-controlled oxidation of bulk SiC. Sebire-Lhermitte et al. (26) identified the presence and location of SiO2 formation in SiC–SiC composites using transmission electron microscopy. They noted the presence of 15-nm-thick SiO2 layers at both the fiber–carbon and matrix–carbon interfaces following an exposure of 1 h at 1123 K in air. That Windisch et al. (20) only observed a weight loss while others (22– 24) observed a weight gain following the weight loss could be the result of lower O2 pressures, exposure time, and perhaps, carbon-layer thickness. The lower O2 pressure would reduce the SiO2 formation rate and therefore the chance for a measurable weight gain during the 5-h exposures. Tortorelli et al. (22,23) used exposures of up to 150 h, whereas Unal et al. (24) used 50-h exposures. Measurable SiO2 formation at pO2 ⬍ 2 ⫻ 103 Pa would need a much greater exposure time than that used by Windisch et al. (20) and even greater than the time used by Tortorelli and More (22). The existence of subcritical crack growth, as described in the next section, that coincides with only weight loss or interphase removal without the embrittling effect of SiO2 or other solid-product formation is the primary difference between an interphase removal mechanism (IRM) and oxidation embrittlement mechanism (OEM) of crack growth. The oxidation results of Windisch et al. (20) demonstrate that the results of Henager et al. (27,28) at temperatures ranging from 1073 to 1473 K and pO2 ⬍ 2 ⫻ 103 Pa occurred by IRM only. The OEM is dependent on SiO2 formation, which depends on O2 pressure, temperature, and time. The interfacial layer thickness may also impact this regime if a solid product can seal off the interface from further reaction. A clear demonstration of this possibility has not been presented but remains as an open issue needing further evaluation. The effect of oxygen on the subcritical crack growth velocity of SiC–SiC is clearly demonstrated by the data given in Fig. 9. Oxygen has little effect on the midpoint displacement (i.e., crack velocity) for about 2 ⫻ 104 s, but a marked increase in the crack velocity is noted for longer times. These tests were performed in the O2 pressure, temperature, and time regime where only weight loss was observed during oxidation studies (20). Therefore, the embrittling effect of 404 Jones et al. Fig. 9 Midpoint displacement of a single-edge notch-beam specimen of SiC–SiC with a C interphase in gettered argon and 3.1 ⫻ 102Pa of O2 ⫹ Ar at 1200°C. a solid reaction product should not be a factor, but only the effect of fiber creep and interfacial removal should have contributed to the crack growth rate. However, even if SiO2 or other glassy phases were present, they would have low viscosity at very high temperatures and would not likely affect the crack growth behavior or cause brittle fracture. The dependence of the crack velocity on O2 partial pressure up to 10% that of air is given in Fig. 10 for tests at 1373 K. There is a sharp increase in the crack velocity at low pressure and a slower increase from pressures of about 0.25 ⫻ 102 Pa up to 2.5 ⫻ 103 Pa. Material with a BN interface exhibited about a factor-of-10 slower crack velocity. Some glass-phase formation was noted in this material as a result of these exposures, but there was no evidence that the crack growth behavior was affected by the presence of this glass phase. The IRM has been observed to cause crack growth in SiC–SiC at temperatures of 1073–1473 K at pressures ranging from 2 ⫻ 102 to 2 ⫻ 103 Pa for stressed samples (29,30). Weight-loss measurements suggest that the IRM operates over these same temperatures and pressures and at 1373 K and a pressure of 2 ⫻ 104 Pa (20). An example of the crack velocity versus crack length for tests on specimens with Hi-Nicalon-reinforced material tested in gettered Ar and Ar ⫹ 2 ⫻ 102 Pa of O2 is given in Fig. 11. The acceleration in the crack velocity induced by the presence of O2 is clearly shown by these data, whereas the effect of temperature on the Ar ⫹ O2 test is only apparent after a crack extension of 3.5 mm. Ceramic Matrix Composites 405 Fig. 10 Minimum crack velocity versus O2 partial pressure for SiC–SiC with a C or BN interphase exposed at 1100°C. Fig. 11 Crack velocity versus crack length for SiC–SiC with a C interphase exposed to 202 Pa of O2 ⫹ Ar (dashed curves) or gettered Ar (solid curves) at 1175°C (1448 K) and 1200°C (1473 K). 406 Jones et al. Dynamic (sample stressed during exposure) OEM has been observed (31– 33) to occur in SiC–SiC at temperatures up to 1073 K, and 1223 K in air (O2 pressure of 2 ⫻ 104), whereas static (sample unstressed during exposure) OEM was also observed in room-temperature tests following elevated-temperature exposures at 1223 K in air (23) and up to 1673 K in dry O2 at a pressure of 105 Pa (24). Heredia et al. (31) reported an upper pest temperature of 1073 K for SiC–SiC tested at elevated temperature in air, whereas Lin and Becher (32) and Raghuraman et al. (33) observed OEM in air at 1223 K. Tortorelli et al. (23) and Unal et al. (24) found the OEM to operate in room-temperature tests following elevated-temperature exposure tests conducted in air without the application of stress. In summary, the upper temperature limit for the dynamic operation of OEM is between 1073 and 1223 K, and the formation of a glass phase at temperatures greater than 1223 K can cause OEM to occur when specimens are tested at lower temperatures. Because the OEM results from the formation of a brittle glass phase, this mechanism must depend on the growth rate and viscosity of the glass phase. The growth rate will increase with increasing temperature and pO2, but the viscosity decreases with increasing temperature. Therefore, there must be a temperature at which the effectiveness of OEM is maximum. The results of Heredia et al. (31), Lin and Becher (32), and Raghuraman et al. (33) appear to define the upper temperature and O2 pressure limits for OEM in SiC–SiC at 1073–1223 K and O2 pressures of 2 ⫻ 104 Pa and above. However, the IRM appears to operate over temperatures of at least 1073–1473 K at O2 pressures of 2 ⫻ 103 Pa and below. It may also operate at temperatures below 1073 K, within the OEM range, at low pressures, but this has not been observed because the crack velocities are too low for experimental measurements. Bouchetou et al. (34) and Frety and Boussuge (35) observed a degradation of the mechanical properties of SiC–SiC containing cracks produced by stress or thermal gradients and exposed to an oxidizing atmosphere at 773 K. The authors did not provide sufficient details to identify the strength loss as OEM; however, OEM is expected at temperatures below about 1223 K, although 773 K would be the lowest reported occurrence to OEM. The low activation for carbon oxidation (50 kJ/mol), as reported by Windisch et al. (20), would result in a small decrease in the oxidation rate of a carbon interphase with decreasing temperature. The transition from OEM to IRM is not only a function of temperature and oxygen pressure but also of the thickness of the carbon interphase layer between the fibers and matrix. Filipuzzi et al. (36) and Cawley (37) noted that interphases with a thin carbon layer (e.g., 0.1 µm) were quickly sealed by SiO2, whereas the carbon was totally oxidized before glass formation in material with a thicker interphase region (e.g. 1 µm). The interphase thickness would not alter the temperature dependence of the OEM to IRM transition but would lower the pO2 and shorten the time for OEM to occur in favor of IRM. A schematic of the matrix, interphase, and fiber oxidation process shown in Fig. 12 from Cawley (37) shows Ceramic Matrix Composites 407 Fig. 12 Schematic of O2 reaction with the C interphase and interphase recession and with the SiC matrix and fiber to form SiO2. the competition between interphase oxidation and SiO2 formation on the matrix and fiber that leads to sealing off of the interphase region. Huger et al. (38) measured the oxidation rates for Nicalon NLM202 fibers exposed to air at temperatures ranging from 700°C to 1200°C. After 100 h, the weight change was about five times greater at 1200°C relative to 700°C with a 1-µm-thick glass layer forming after 100 h at 1000°C. Clearly, the oxide layer can grow sufficiently thick to seal the interphase region and to perhaps to act as a crack initiator which could reduce the fiber strength. The environmental stability of composites with silicon-based matrices, such as blackglass, nitrides, and carbides have also been evaluated by several authors. In studies of material with Nextel 312 fibers in a matrix of Allied Signal Blackglass with a BN interphase, Campbell et al. (39) measured the weight change and bend strength of material exposed to dry air, air ⫹3% H2O, dry air ⫹ 80 ppm KCl vapor, and air containing 3% H2O ⫹ 80 ppm KCl vapor. The exposures were conducted for variable lengths of time at 700°C and 5 h at 900°C. The composites show only continuous weight loss with time for exposures to dry air at 700°C, but a weight gain following a small initial weight loss in KCl containing environments. The weight gain was identified as resulting from the formation of alkali silicates. Vaidyanathan et al. (40) also studied the effects of 408 Jones et al. oxidation on the mechanical properties of the Nextel 312/BN/blackglass composite but for times of 20–1000 h at 600°C. They observed a 50% reduction in strength after 500 h at 600°C and concluded that times greater than 200 h under these conditions were a concern for the durability of the composite. Oxidation resulted in increased fiber pullout, consistent with IRM, but at a much lower temperature than IRM in a SiC–C–SiC composite. The authors concluded that the degradation may have resulted from the weakening of the fiber–matrix interface. The instability of the fiber–matrix interphase in SiC–SiC composites has led to the evaluation of coatings for these materials. Fox (41) evaluated the oxidation resistance of three coatings on SiC–SiC composites: (a) CVD SiC, (b) a particulate-based sealant with a CVD SiC outer layer, and (c) a boron-rich inner layer and CVD SiC outer layer. Oxidation studies were conducted in dry oxygen at 981–1316°C. All three provided protection for up to 100 h, and the CVD SiC was the most protective. Oxygen diffusion through the SiO2 that formed on CVD SiC was considerably slower than through the glass layer that formed on the other two coatings which contained B. Lee and Miller (42) evaluated the stability of mullite coatings with a refractory oxide barrier coating on SiC–SiC exposed to air at 1200–1400°C. The sample temperatures were cycled every 1, 2, or 20 h from 1200°C and 1300°C to room temperature and every 1 h from 1400°C to room temperature. The mullite–refractory oxide composite coating exhibited improved adherence and oxidation resistance relative to a straight mullite coating. D. Hydrogen-Containing Environments Herbell et al. (43) have evaluated the thermodynamic stability of SiC in pure H2 at 1 atm, as shown in Fig. 13. The primary gaseous reaction produce is CH4 as described by Eq. (2), whereas other reactions which produce SiH4 and SiH are also possible at temperatures as low as 900°C. A small amount of H2O can alter the phase stability such that at 1400°C and about 1000 ppm of H2O, the dominant gaseous reaction products become SiO and CO (Fig. 14). Results for lower temperatures were not reported, but the reaction of H2O with SiC occurs at much lower temperatures, so similar reaction products would be expected at lower temperatures. No loss in the room-temperature flexural strength of sintered SiC was noted by Herbell et al. (43) for samples exposed to H2 saturated with H2O for 100 h at temperatures from 800°C to 1400°C. In dry H2 (25 ppm H2O), Hallum and Herbell (44) noted a 33% decrease in the fracture strength of sintered SiC after exposures of 500 h at 1100°C and 1300°C. A statistically significant decrease in flexural strength was also observed after 50 h at 1000°C. The stability of SiC in an Ar–H2O–5% H2 mixture was calculated by Jacobson et al. (45) in the same manner as the H2 –H2O mixtures, with the result shown in Fig. 15. Except for the lower gas pressures and the shift in the relative activities of SiO and CH4 in region III, the results are essentially Ceramic Matrix Composites 409 Fig. 13 Gases for equilibrium partial pressures of reaction products for reaction of SiC with pure H2 at 1 atm. identical. At 1300°C, Jacobson et al. (45) measured a weight loss of around 1 mg/cm2 after 24 h in a region II Ar–H2 –H2O mixture. Hydrogen may also react with the carbon interphase to form CH compounds. This reaction would be in addition to the direct reaction with matrix and fiber as described by Herbel et al. (43). Springer et al. (46) evaluated the reaction of Ar ⫹ H2 environments on the weight change of SiC–SiC composites which had a carbon fiber–matrix interphase. They used a thermogravimetric analyzer Fig. 14 1400°C stability of SiC in H2 ⫹ H2O at 1 atm. 410 Jones et al. Fig. 15 Thermodynamic analysis of SiC ⫹ 5% H2 –Ar at 1300°C. All pressures are in atmospheres. to study the weight change at 1000–1200°C with Ar ⫹ 0.1% H2 and Ar ⫹ 1.0% H2. They found that the reactivity of the carbon interphase to H2 was substantially less than to O2. For instance, Ar ⫹ 100 ppm O2 produced a weight loss 26 times greater at 1000°C relative to Ar ⫹ 0.1% H2. Nightingale (47) found an activation energy of 65 kcal/mol for H2 reacting with bulk graphite, whereas Springer et al. (46) found activation energies of 18 and 34 kcal/mol for 0.1% and 1.0% H2, respectively. The carbon interphase material is a mixture of amorphous carbon and graphite, so that the lower activation observed for the carbon interphase material could be the result of the lower stability of the interphase material relative to bulk graphite. A conclusion of the study by Springer et al. (46) is that H2 is much less of a concern than O2, but that for environments with low pO2, the reaction of both SiC and C with H2 could be a significant environmental stability issue, especially at temperatures above 1200°C. III. OXIDE MATRIX COMPOSITES The chemical instability of the SiC in the presence of alkali elements and the fiber–matrix interphase in SiC–SiC composites in oxidizing environments are Ceramic Matrix Composites 411 factors that encourage the development of oxide matrix composites. Much of the high-temperature corrosion data for oxide matrix composites exists for particulate-, whisker-, or platelet-reinforced material, which have not been optimized for strength and toughness. Oxide fiber development has progressed to the state where continuous-fiber composites are being produced; however, there is no high-temperature oxidation data for these materials. Examples are alumina and alumina–YAG matrix composites reinforced with Nextel 610 and 720, as reported by Goettler (48). Interphase layers of ErTaO4 or CaWO4 are being evaluated for producing fiber pullout and fracture resistance. However, the stability of these interphase materials in oxidizing, reducing, or salt environments has not been evaluated. Even though the matrix and fiber may exhibit excellent behavior in oxidizing environments, uncertainty about the composite chemical stability remains. Also, the high-temperature strength of oxide fibers is less than SiC fibers, such that further improvements in strength must occur before continuous-fiber oxide–oxide composites are attractive for high-temperature applications. Borom et al. (49) have examined the oxidation behavior of Al2O3 reinforced with SiC and MoSi2 particles and SiC whiskers. The particulate volume fractions ranged from 10% to 30% and tests were conducted at 1200–1500°C in air; the oxidation rate was determined by weight change and reaction layer thickness. Both SiC and MoSi2 form protective SiO2 layers when oxidized as bulk materials. Borom et al. (49) reported a 15-fold increase in the oxidation rate of these phases when incorporated into an Al2O3 matrix. This increase was postulated as resulting from the volume change of the reaction product that forms on the composite and the thermal expansion mismatch of the reaction product with the composite. Both of these factors were less favorable for the composite as compared to bulk SiC and MoSi2. Larger volume fractions of these phases produced a large volume fraction of mullite in the reaction scale and this was favorable because the silica in the mullite will produce a more viscous scale that will allow more stress relaxation and accommodation for mismatch stresses. These authors suggested that a mullite matrix is preferred because the reaction product will contain aluminosilicate plus mullite, which will flow and relax thermal mismatch stresses. The bend strength of SiC-whisker-reinforced (28 vol%) Al2O3 was found by Leaskey et al. (50) to increase by 33% when oxidized in air at 1600°C for 15 min. Composites with SiC particle reinforcement showed a 66% improvement in the bend strength following an oxidizing treatment of 2 h at 1600°C. The authors suggested that the improved properties are the result of the oxidation of the SiC reinforcement to produce a compressive surface layer. The following conditions were necessary for this improvement: (a) a sufficiently large SiC content to produce a continuous oxide surface layer, (b) oxidation conditions that produce a low porosity layer with a critical thickness, and (c) elimination of large flaws in the bulk of the material. 412 Jones et al. The tensile strength of Nicalon-fiber-reinforced Al2O3 following heat treatment in air at 750°C has been reported by Heredia et al. (31). This material contained 0°/90° fiber orientation with a BN–SiC interphase. The room-temperature tensile strength was found to decrease from about 250 MPa to about 120 MPa following exposure to air at 750°C for 24 h. The formation of a glass phase on the Nicalon fiber was suggested as the cause of the observed oxidation embrittlement. Corrosion studies of SiC-reinforced Al2O3 have been conducted in coal combustion environments by Watne et al. (51) and Breder et al. (52). In the study reported by Watne et al. (51), the material produced by the Lanxide Corp. contained 50% SiC with 10% residual Si. Following a 100-h pilot-scale combustion test at about 1350°C in a radiant zone of the furnace, the composite, in the form of a tube, was intact but had a 0.85-mm reduction in the wall thickness. This loss was suggested as resulting from erosion from slag flow on the tube. The original wall thickness was 5.25 mm. A smaller amount of loss was found for a tube placed in the convective pass region of the combustor where the temperature was about 1200°C. Breder et al. (52) exposed a similar tube made by Lanxide Corp. to coal slag obtained from two coal combustion plants. Exposures were conducted in a box furnace with the tube and coal slag at temperatures of 1090°C, 1260°C, and 1430°C. Fracture tests were conducted on samples removed from the tubes following a 500-h exposure. The tube strengths were reduced by 20–45% at 1260°C depending on the type of slag. Although Al2O3 is the most commonly used matrix for oxide matrix composites, composites with other oxides such as MgO, ZrO2, and mullite have also been evaluated. The oxidation kinetics of SiC particulate-reinforced MgO has been examined by Hallum (53) and Camey and Readey (54). Hallum studied MgO reinforced with 5, 10, and 15 vol% SiC particles or whiskers over the temperature range of 1100–1500°C. The reaction-product thickness increased with the square root of time and was a function of the volume fraction of SiC in the composite. Mg cation diffusion was proposed as controlling the growth rate with a reaction layer formed by Mg cation diffusion through the reaction layer to the atmosphere where oxidation produced a columnar growth region. Camey and Readey (53) identified three oxidation-product layers unlike the single layer observed by Hallum; however, they agreed with Hallum regarding the growth rate being controlled by Mg cation diffusion through the product layer. Luthra and Park (55) evaluated the oxidation of SiC in mullite and alumina matrices and found parabolic rate constants that were three orders of magnitude larger than SiC. Xu et al. (56) measured the effects of adding ZrO2 to mullite on the oxidation of mullite–zironia–SiC composites. They found that the addition of ZrO2 to mullite–SiC composites increased the reaction rate with oxygen. They rationalized this as being due to the increased diffusion rate of oxygen in the zirconia phase. A rapid ‘‘mode II’’ type of oxidation, where oxygen can penetrate Ceramic Matrix Composites 413 deep into the sample before the outer region is completely oxidized, occurred at 1200–1400°C and with the volume percent of ZrO2 greater than 20%. IV. GLASS MATRIX COMPOSITES Glass and glass–ceramic matrix composites are the most developed class of ceramic matrix composites. These composites are easier to prepare than SiC–SiC or oxide matrix composites and so have received further development and evaluation than other CMCs. The glass matrices employed in these composites include calcium–aluminosilicate (CAS), lithium–aluminosilicate (LAS), magnesium– aluminosilicate (MAS), and barium–magnesium–aluminosilicate (BMAS). There have been a number of microstructural, mechanical property, and environmental effects studies of materials reinforced with Nicalon-type fibers. A. High-Temperature Air Environments Alteration of the fiber–matrix interface is one of the primary effects of oxidation on glass matrix–Nicalon composites. Daniel et al. (57) evaluated the oxidation of CAS–Nicalon composites over the temperature range of 375–600°C in air for 100 h. They evaluated the change in the fiber–matrix interfacial properties with a nanoindentation push-down test and four-point bend tests. At exposure temperatures of 450°C and above, the composites exhibited brittle failure with minimal fiber pullout. The transition from tough behavior with fiber pullout for lowertemperature exposures to brittle fracture was associated with an increase in the fiber–matrix frictional shear stress. This increase in the frictional shear stress is accompanied by the loss of the fiber–matrix interfacial carbon layer and the resulting residual stress causing the matrix to apply a compressive stress to the fiber. This clamping stress on the fiber reduces fiber pullout and causes brittletype behavior. Microstructural evaluation of the fiber–matrix interfacial region of CAS and LAS–Nicalon fibers exposed to air at 600°C or 900°C have been reported by Cooper and Chyung (58). The oxidized foils were very fragile, consistent with the embrittlement noted by Daniel et al. (57). The interface was found to have a silicate composition instead of the graphite composition. This is in contrast to the conclusion reached by Daniel et al. (57) that the loss of the graphite layer by oxidation resulted in a clamping stress on the fiber and the resulting brittle‘‘type’’ fracture. The formation of a silicate that forms a strong bond between the fiber and matrix will accomplish the same result and would also be consistent with the increased interfacial shear stress observed by Daniel et al. (57). High-temperature mechanical property tests of BMAS–Nicalon composites in air by Sun et al. (59) showed only limited oxidation of near-surface fibers in 414 Jones et al. tests where the stress was below the proportional limit. However, dynamically loaded samples loaded above the proportional limit, or matrix cracking stress, exhibited limited fiber–matrix interface oxidation. Oxygen diffusion along matrix microcracks created by stresses exceeding the proportional limit was thought to be the primary cause for the fiber–matrix interfacial oxidation. This effect was most pronounced under cyclic loading compared to static or quasistatic loading. Embrittlement of a MAS–Nicalon composite during fatigue loading at 500°C in air was also reported by Heredia et al. (31). A reduction in fatigue life was noted after only 1000 cycles at 500°C relative to room-temperature tests. Heredia et al. (31) related this loss to a ‘‘pest’’ process where the Nicalon is embrittled and they suggest that the compressive matrix stress in the glass-ceramic matrix– Nicalon composites requires a cyclic stress to reveal this process. Sorensen et al. (60) studied the effect of environment and frequency on the fatigue properties of a CAS–Nicalon composite. They concluded that fiber–interfacial wear processes play a significant role in the loss of fatigue life of these composites and that the environment enhances this wear-induced loss of strength. B. Hot Corrosion Environments High-temperature salt environments will occur in engine components on a Navy gas turbine engine and heat exchangers in coal-fired power plants. There is a strong emphasis on increasing the trust-to-weight ratio of Navy planes, and the low-density and high-temperature performance of CMCs are needed to achieve these goals. CAS–Nicalon and LAS–Nicalon composites have been evaluated for this application by Wang et al. (61,62). They examined the reaction of sodium sulfate with these composites by coating specimens and heating them to 900°C in either air or argon atmospheres for up to 100 h. The CAS–Nicalon composites exposed in air showed surface cracking and extensive reaction between the salt and the Nicalon fibers. The surface fibers were completely attacked and were totally removed. X-ray diffraction was used to identify the presence of CaSiO3 and NaAlSiO4. The unreinforced CAS glass exposed to the same conditions reacted to form NaAlSiO4 but not CaSiO3. Therefore, the SiC fibers contributed to the reaction products and altered the corrosion reaction. The tensile strength and strain to failure of the CAS–Nicalon composite exposed to sodium sulfate in air was reduced relative to the as-received properties and those for material annealed at 900°C for 100 h but without the presence of the salt. However, the properties of material exposed to salt in an argon atmosphere showed no degradation in properties. The authors concluded that oxidation is the primary reaction responsible for the strength degradation of the composite. In contrast, the LAS–Nicalon composites did not form additional phases, although there were surface cracks and interdiffusion of Na into the composite and Mg outward diffusion. A 30% Ceramic Matrix Composites 415 strength reduction was noted for the LAS–Nicalon composite exposed to the salt, presumably a result of the surface cracking and Na and Mg interdiffusion. A thermodynamic evaluation was conducted by Kowalik et al. (62) for the CAS–Nicalon composite exposed to sodium sulfate at 900°C, as reported by Wang et al. (61). This study suggested the following reaction path: (1) SiC oxidizes to form SiO2, (2) the silica reacts with the Na2O in Na2SO4(Na2O⋅SO3), (3) the result of reaction 2 may lead to a liquid oxide (or soda slag) phase which may attack the CAS matrix, (4) the SO3 from reaction 2 may combine with the CaO in the matrix to form CaSO4, and (5) this lost CaO from the matrix is replaced by Na2O to yield NaAlSi3O8. These thermodynamic predictions closely match the experimental results reported by Wang et al. (61). Step 1 shows the significance of oxidation in the high-temperature corrosion of these materials. High-temperature corrosion studies of CAS–Nicalon and BMAS–Tyranno exposed to sodium and magnesium salts have also been conducted by Scott et al. (63). The environments were 3.5% NaCl, 3.5% magnesium salts, and a mix of 3.5% of both sodium and magnesium salts. The samples were coated with these solutions by immersion and then heated to 600°C, 800°C, or 1100°C for up to 60 h. Reaction occurred primarily between the Ca and Mg ions, and the Nicalon in the CAS–Nicalon composite, but the Na ions penetrated the glassy phase and lowered its viscosity in the BMAS–Tyranno composites. The authors concluded that both reactions were of concern for the stability of these composites in high-temperature salt environments. V. SUMMARY Ceramic composites are being considered for a variety of high-temperature applications in which their corrosion properties will be important for their performance. Examples include combustor liners and blade shrouds for gas turbines, heat exchangers in a coal-fired power plants, burner nozzles, gas injection lances, sensor shields, tundish nozzles for molten Al and steel plants, and furnace/reformer tubes. Each of these applications involves some form of corrosion. The corrosion of ceramic composites is more complicated than that of a monolithic ceramic because composites are chemically and microstructurally heterogeneous. The high-temperature corrosion of CMCs are often affected by the fiber, fiber–matrix interphase, or the method used to produce the matrix of the composite. High-temperature oxidation of the C or BN interphase in SiC–SiC composites is a clear example where the interphase causes the corrosion performance of the composite to be less than that of monolithic SiC. The presence of the SiC in mullite or alumina matrix composites were also found to increase the parabolic rate constants for oxidation by several orders of magnitude, whereas 416 Jones et al. the presence of the SiC fiber resulted in a different reaction product in a CAS– Nicalon composite than in the unreinforced matrix when reacted with a hightemperature salt environment. 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In: R Naslain, J Lamon, D Dougmeingts, eds. Proceedings of High-Temperature Ceramic-Matrix Composites—I. Abbington, Cambridge: Woodhead Publishing, Ltd., 1993, p. 343. 61. S-W Wang, RW Kowalik, R Sands. Ceram Eng Sci Proc 385, July–Aug. 1993. 62. RW Kowalik, S-W Wang, PD Ownby, DM Thompson, WT Thompson. Ceram Engineer Sci Proc 893, Sept.–Oct. 1995. 63. V Scott, S Bleay, R Cooke. In: R Naslain, J Lamon, D Dougmeingts, eds. Proceedings of High-Temperature Ceramic–Matrix Composites—I. Abbington, Cambridge: Woodhead Publishing Ltd., 691, 1993. 14 Issues in Predicting Long-Term Environmental Degradation of Fiber-Reinforced Plastics Aaron Barkatt The Catholic University of America, Washington, D.C. I. INTRODUCTION Several major unresolved issues are involved in predicting the effects of environmental degradation on the long-term behavior of fiber-reinforced plastics in construction applications. Many types of change in mechanism in the course of exposure of fiber-reinforced plastics (FRPs) to the surrounding environment are possible. Such changes limit the applicability of extrapolation from short-term test data and, in certain cases, cause the degradation rate to rise in the course of the exposure. Such increases in degradation rate may be gradual or abrupt, and their effects may rapidly disappear or persist for long periods of time. Changes in the degradation mechanism may involve the degradation of the fibers, the matrix, or the interphase. Changes in mechanism also affect the dependence of the degradation rate on temperature and moisture content and, thus, limit the range of conditions over which temperature may be used as an accelerating factor in predictive tests. In addition to these scientific issues, uncertainties concerning the effects of scale, limitations of existing test procedures, and, in particular, the variability of FRPs produced on a large scale and the lack of information regarding the effects of such variability on the chemical properties of the materials constitute problems that require solution as a prerequisite for extensive use of FRPs in large-scale structures, such as the bridges, highways, and buildings. This chapter does not attempt to present an updated, comprehensive overview of the area of environmental degradation of fiber-reinforced plastics. In419 420 Barkatt stead, it attempts to use published data to focus on some of the serious issues that complicate the development of long-term predictions of the effects of the environment on the behavior of FRPs in construction applications. These issues include scientific complications involving changes in degradation mechanism, in particular changes that make it possible for supralinear behavior to take place, as well as engineering problems such as scaling, test design, and, in particular, the variability of existing materials and the lack of information about the effects of such variability on the reliability of test-derived predictions of material performance. II. FIBER-REINFORCED PLASTICS AS CANDIDATE MATERIALS FOR CONSTRUCTION APPLICATIONS As the use of FRPs in various applications has increased, studies have been carried out on the environmental degradation of FRPs and its effect on performance. The results of many of these studies are summarized in several review articles and compilations, such as those listed in Refs. 1–6. The present chapter is not intended to be a comprehensive review of environmental effects on FRPs. Rather, it attempts to highlight some of the issues and considerations involved in expanding the use of FRPs to high-volume, low-cost applications in civil construction. The development of FRPs started during World War II as a result of the search for lightweight materials possessing high strength and high stiffness for aircraft structures and rocket motors. Subsequently, the use of FRPs in the aerospace industry has expanded, and these materials found, in addition, a broad range of other applications. The development of advanced molding techniques, for instance, has opened the way for the use of FRPs in the automotive industry, and they have also been extensively used in other transportation systems, such as light rail and marine vessels. FRPs are used in a large number of specialty products, including electric and electronic equipment and consumer products (e.g., sports and recreational equipment). In the cases of most of the applications mentioned so far, performance, rather than cost, has been the major consideration. As a result, it has been possible to use relatively expensive ingredients, such as carbon or boron fibers, in such applications along with sophisticated techniques of fabrication and of quality control. The good results obtained with FRPs have led to the expansion of their use to applications on a larger scale, including light industrial structures such as stairs, platforms, and rails, as well as chemical storage tanks, liners, and reaction vessels. FRPs have also been introduced into use in pipelines and oil storage tanks by the petroleum and petrochemical industries. This, in turn, has stimulated interest in the use of FRPs in large civil structures such as supports for bridges, roads, and buildings. However, large-scale use of Long-Term Degradation of FRPs 421 FRPs in such structures requires addressing issues that include (a) cost as a major factor, (b) a very broad range of environments, often subject to numerous and frequent changes, and often involving the presence of highly corrosive media such as water previously in contact with cement, and (c) the need to maintain the performance of the FRP-supported structures not far below initial levels over long periods of time, covering several decades. For instance, service life of 75 years is required for highway structures (7). The combination of these factors requires extensive work in the areas of testing and predictive modeling to guide efforts in FRP fabrication, quality control, and structural design in order to ensure adequate safeguards against premature degradation of performance under service conditions. A. Structure of FRPs Fiber-reinforced plastics consist of a combination of fibers, which impart strength and stiffness to the material, with a polymeric matrix, that serves to hold the fibers in the desired configuration (random or aligned), to transfer applied stresses to the fibers, and to protect the fibers against abrasion and corrosion. The boundary region between the fibers and the matrix is the interphase, consisting of the bonding layer interfacing with both the fiber and the matrix (8). This region is of critical importance in controlling the adhesion of the fibers to the matrix and the ability of the composite as a whole to maintain its strength in moisture-containing environments. The adhesion of the fibers to the matrix results from a combination of mechanical fitting and chemical adhesion. The relative importance of the physical and chemical factors, respectively, to fiber–matrix adhesion, as measured, for instance, in fiber pullout tests, is variable and is not yet fully understood in many cases. Empirically, however, it is well established that the use of silane coupling agents to treat the fibers, known as sizing, improves the properties of the interface. Sizing just prior to impregnation of the fibers with resin improves the adhesion between the inorganic fiber and the organic resin and enhances the resistance to moisture. [Such permanent sizing should be distinguished from pretreatment of the fibers with organic lubricants to minimize fiber–fiber abrasion and to facilitate maintaining the fibers in their desired configuration. These organic lubricants are usually removed by heating before the treatment with the permanent sizing agent (i.e., the silane coupling agent)]. In addition to the treated fibers and the matrix, FRPs often contain additional components such as fillers (typically clay or hydrated alumina) and mold-release agents (lubricants). Polymeric matrices can be divided into two broad classes consisting of thermosetting polymers and thermoplastic resins, respectively. Thermoplastic resins such as polyetherimide (PEI), polyethersulfone (PES), polysulfone (PSU), and polyetheretherketone (PEEK) generally possess higher tensile strength, better 422 Barkatt impact resistance, and much higher heat-distortion temperatures than thermoplastic polymers. However, the cost of thermoplastic matrices is much higher than that of thermosetting matrices (9). Accordingly, thermosetting matrices have a greater potential at the present time for use in civil construction applications, whereas thermoplastic matrices are more suitable for high-performance engineering applications such as aerospace components (8), especially in combination with high-performance carbon fibers. Major thermosetting matrices include vinylester, epoxy, polyester, and phenolic resins. Vinylester and epoxy matrices are somewhat more expensive than the other two types of polymers, but they possess higher tensile strength as well as better resistance to attack by water and by various chemical solutions (9). Of the four types of thermosetting resins mentioned, phenolic resins are the least expensive, but their relatively low tensile strength and their susceptibility to void formation during curing, which results in large capacity for absorption of moisture (10) render them less promising for civil structural applications than the other three types of polymeric matrix material. Nevertheless, phenolic composites are suitable for a variety of applications in consumer products, electrical equipment, and automotive components. Other, more advanced types of thermosetting resins, developed for high-temperature applications, consist of polyimides and polybismaleimides (BMIs). Because of their brittleness, high cost, and the requirement for sophisticated techniques during fiber incorporation (10), the use of polyimides and BMIs is largely confined to the aerospace industry. Accordingly, the most promising candidates for use in large-scale civil applications are polyesters, epoxies, and vinylesters. Vinylesters combine the desirable thermal, mechanical, and chemical characteristics of epoxies with the rapid curing and ease of processing of polyesters (11). The three major types of fiber used in the fabrication of FRPs are glass fibers, carbon fibers, and aramid fibers, respectively. Carbon fibers exhibit the highest strength and stiffness. They have excellent chemical resistance as well as high-temperature performance, which makes it possible to use these fibers in conjunction with carbon, metal, or ceramic matrices in demanding aerospace applications. In civil construction applications, however, such high-temperature resistance is not required, and the cost of carbon fibers poses major hindrance to their use. In addition, carbon fibers are moderately brittle, resulting in a low capacity to absorb impact energy. They have low abrasion resistance and are subject to galvanic corrosion in aqueous media in the presence of metals and alloys (12–14). Aramid fibers, produced from poly(paraphenylene terephthalamide) have lower strength than carbon or glass fibers, but they can absorb large impact energies. Their transverse strength and compressive strength are relatively low and they exhibit a strong tendency to absorb water. In addition, their adhesion to polymeric matrices is sometimes weak and composites based on such fibers are difficult to cut and machine. Glass fibers have high tensile strength, although not as high as that of carbon fibers. They have a high capacity to absorb impact Long-Term Degradation of FRPs 423 energy and a good temperature resistance. Unlike aramid fibers, glass fibers do not tend to absorb water and they have excellent temperature resistance. Glass fibers are less stiff than carbon fibers. The low cost of glass fibers is a major consideration that makes them prime candidates for use as reinforcements in composites intended for large-scale construction applications. In addition to the three major types of fiber described, advanced fibers, such as boron fibers and metal fibers, have found use in the aerospace industry and in certain sporting goods, but the excellent temperature resistance that such fibers possess is usually unnecessary when used in conjunction with polymeric matrices, and their excellent mechanical and chemical stability is offset by their high cost in large-scale applications. Polyethylene fibers have high tensile strength and stiffness, but their compressive strength and modulus are low and they have very low softening temperatures. B. Environments The environments that FRPs are expected to encounter in civil construction applications can be divided into those expected in everyday use and those encountered only under accident conditions. Routine service environments involve a range of temperatures between about ⫺30°C and 60°C, depending on geographical location and subject to seasonal and diurnal fluctuations. The amount of moisture in the environment may vary between a relative humidity of approximately 30% and fully water-saturated air. Contact with liquid water will occur as a result of rain, melting snow, or runoff. The water may be slightly acidic (ordinarily with a pH no lower than 4.5–5) in the case of acidic precipitation or extraction of acidic components from the ground. On the other hand, very high pH values may be encountered in cases where FRP tendons are used in direct contact with cement environments, as the pH of cement pore water is as high as 13–14 (15). The salt content of the water may range from very low in the case of fresh rainwater to very high in the case of slush containing large amounts of road salts (NaCl or CaCl2). Both wet–dry cycling and freeze–thaw cycling may aggravate the effects of aqueous environments. FRPs in civil engineering applications can be expected to have some exposure to gasoline and motor oil fumes, and occasionally to gasoline and oil in liquid form, but it has been found that such nonpolar media have smaller effects on the mechanical properties of FRPs than aqueous media (16). Some exposure to solar radiation in the ultraviolet (UV) and visible ranges may take place, but at levels much lower than those encountered, for instance, in sails and boat surfaces. While exposed to the environments described earlier, FRP reinforcements in civil structures such as bridges, roads, and buildings are subject, of course, to large and variable mechanical stresses. Although discussion of the effects of such stresses by themselves is outside the scope of this review, the interaction between environmental and mechanical stresses has to be taken 424 Barkatt into account in situations involving phenomena such as fatigue and stress corrosion. III. PREDICTION OF THE LONG-TERM ENVIRONMENTAL STABILITY OF MATERIALS A. Accelerated Tests and Service-Condition Tests The general approach for predicting the long-term performance of structural materials consists of a combination of tests and modeling efforts (17). This approach is based on the use of accelerated tests as well as service-condition tests. Servicecondition tests, performed over the range of conditions to which the material is expected to be exposed in service, are, of course, extremely useful, especially when they are conducted under the most aggressive conditions that the material may experience. However, the duration of such tests is limited, due to practical considerations, to periods much shorter than the required service life. Accordingly, it is necessary to supplement such tests with accelerated tests, carried out under conditions more severe than those expected under the most extreme service environments. An acceleration factor F, defined as the ratio k a /k s between the degradation rates in the accelerated test and under service conditions, respectively, is experimentally determined. Based on this ratio and on the experimentally determined time it takes the material to fail under the accelerated conditions, T a, the service life, T s, can be evaluated from the relationship T s ⫽ T a ⋅F ⫽ T a 冢冣 ka ks (1) Various methods of accelerating tests of environmental degradation have been considered and tried. The principal acceleration methods include raising the temperature, the use of a more corrosive environment, and the application of mechanical stresses during the exposure. Acceleration through the use of elevated temperatures is the most straightforward and most commonly used of these acceleration methods. In many cases, the degradation rates of materials exhibit an Arrhenius-type dependence on the temperature over certain temperature ranges. In such cases, the acceleration factor F can be expressed as F⫽ ka ⫽ exp ks 冤冢ER 冣 (T a ⫺1 s 冥 ⫺ T a⫺1) (2) where T a and T s are the temperatures of the accelerated test and the temperature corresponding to the service condition, respectively, E a is a constant activation energy characteristic of the degradation of the particular material under evalua- Long-Term Degradation of FRPs 425 tion, and R is the gas constant. In the case of silicate glasses, for instance, Arrhenius-type behavior has been observed over a broad range of temperatures (up to at least 90°C), and the values of E a were generally observed to range between 20 and 25 kJ with respect to leaching of alkalis and between 15 and 20 kJ with respect to dissolution of the silicate matrix (18,19). Accelerated tests based on the use of elevated temperatures have been used in the evaluation of FRPs for structural applications in the aircraft industry to predict the effects of moisture as well as of temperature within the range of expected service environments (20). The use of accelerated tests is very useful in predicting the performance of materials over periods of time longer than those available for service-condition testing. However, the use of this approach is limited to the range of conditions over which Eq. (1) is valid. This range is generally limited to those conditions under which the mechanism controlling the kinetics of the degradation of the material is the same as the controlling mechanism under service conditions. The use of accelerating conditions that cause a different degradation mechanism to become predominant generally results in inapplicability of the accelerated test to service conditions. It is known, for instance, that in the case of metals, the structure of the oxide layer and, consequently, the power law characterizing the oxidation kinetics are different at high temperatures than they are at the relatively low temperatures used in many applications (21,22). Likewise, the nature of the alteration products that control the course of glass hydration depends on the temperature (23). Similar precautions (i.e., assuring that the acceleration method does not cause a change in the nature of the controlling degradation mechanism) are required upon the use of other acceleration methods such as employing more corrosive environments or the application of mechanical stresses. B. Degradation Kinetics and Predictive Modeling for FRPs In the case of composite materials, the role of model development in establishing predictions of long-term performance is extremely important. As detailed earlier, in certain cases involving single-phase materials the kinetics of degradation follows a simple, uniform-rate law characterized by a single rate constant that exhibits a simple Arrhenius-type behavior. In such cases, the extent of long-term degradation can be readily correlated with the results of short-term accelerated tests using simple extrapolation. However, such simple behavior is not very common even in the cases of homogeneous materials, and the probability of encountering such behavior in the cases of complex material systems such as composites is very low. The development of more sophisticated predictive models to evaluate long-term performance and the use of test data to verify such models are therefore the key for the use of composites in applications involving long-term exposure in service environments. Considerable efforts have been made to model the degra- 426 Barkatt dation of FRPs. Unfortunately, comprehensive models are not yet available to provide a quantitative basis for evaluating the performance of FRPs as construction materials. The general considerations involved in the development of accelerated tests for FRPs and the use of such tests in predictive modeling have been surveyed (20,24). The starting point for many modeling efforts is broadly viewing the degradation of FRPs in aqueous or humid environments as a sequence of two processes, the first of which is penetration of moisture into the composite material and the second is hydrolytic attack on the structure. The ingress of moisture has been considered by many authors in terms of Fickian diffusion (25). Such models make it possible to characterize the accumulation of moisture in the composite as a function of time until the attainment of a maximum moisture content, at which time the material can be considered to become saturated with respect to its capacity to absorb moisture (25,26). Furthermore, the dependence of the diffusion coefficients on temperature has been considered by many authors in terms of a simple Arrhenius-type dependence, with logD varying linearly with 1/T, where D is the diffusion coefficient and T is the absolute temperature (25,27). This picture yields a complete description of the first stage of the hydrolytic attack in terms of a constant moisture content at saturation and a constant activation energy applicable to the temperature dependence of the diffusion coefficient (28). Unfortunately, the range of applicability of the simple approach described here is limited by the existence of several complicating phenomena. It has been pointed out (24) that the spatial nonuniformity of degradation in FRP materials converts originally simple specimens into complex structures with nonuniform chemical and mechanical states. Overall modeling of the resulting complex degradation behavior requires combination and coupling of submodels of diffusion, degradation, matrix shrinkage, mechanical property loss, and effects at the laminated-plate levels. The limitations of each submodel have to be taken into account. For instance, the validity of submodels based on mass-loss rates obeying Arrhenius-type behavior was found to be mostly limited to comparative evaluations (24). It is a general observation that a particular kinetic law can only be expected to hold as long as a particular process controls the overall rate of degradation. In the case of FRP hydration, deviations from the simple picture described occur when the saturation limit of the moisture content is no longer constant, but increases as the exposure continues as a result of swelling, increase in hydration stress, and the resulting formation of microcracks and microvoids (29). Another limitation, which affects the polymeric matrix (as well as reinforcing fibers made out of polymeric materials such as aramid), is the fact that moderate increases in temperature bring such polymers to the glass transition (T g) range. As the temperature reaches this range, progressively larger regions of the polymeric matrix begin to move and rearrange their orientation with respect to one another Long-Term Degradation of FRPs 427 (30). Exposure to temperatures in excess of T g can facilitate the penetration of moisture and the extent of the resulting hydrolytic attack. On the other hand, such exposure may lead to an increase in the extent of cross-linking, resulting in an enhancement in rigidity and resistance to hydrolysis. This effect is known as postcuring (31). Models have been developed for changes in T g as a result of exposure to humid environments at elevated temperatures (25). A third example of possible changes in mechanism during environmental exposure involves changes in the composition and reactivity of the aqueous phase during contact with the composite material. Such changes may lead to an increase in acidity around the polymeric matrix due to extraction of acidic monomers (e.g., acrylic acid), as well as to local increase in basicity around reinforcing fibers made out of silicate glass as a result of selective leaching of alkali and alkaline earth ingredients of the glass (32). Changes in mechanism due to effects such as the three phenomena discussed here were shown to make it possible for FRP corrosion rates to exhibit supralinear behavior consisting of considerable increases after a certain period of exposure. Such increases were observed in measurements of dissolution rates (see Fig. 1) (29,32) as well as in measurements of mechanical properties. Thus, upon prolonged exposure to humid air, the tensile strength of a sheet-molding compound was found to remain constant for the first 2 months and then to fall off sharply (see Fig. 2) (16). Thus, upon repetitive wet–dry cycling, the amount of damage per cycle was observed to increase with the increasing number of cycles (see Fig. 3) (33). The rate of increase in the magnitude of various indicators of the effect of service conditions, such as strain, was found in certain cases to Fig. 1 Normalized dissolution rates of silica in leachates from exposure of a vinylester/ glass rod to deionized water. (From Ref. 29.) 428 Barkatt Fig. 2 Changes in ultimate tensile strength, short beam shear strength, tensile modulus, shear modulus, and weight of a glass/polyester sheet molding compound in humid air as a function of immersion time: (— — —) ‘‘baseline’’ data; (䊉) specimens tested after 6 months immersion followed by 3 weeks of drying; (䊊) specimens tested without postdrying. Bars indicate spread in data. Left: 23°C; Right: 93°C. (From Ref. 16.) exhibit a distinct rise at a certain time after the beginning of the test. For instance, upon subjecting sheet-molding compounds to a combination of mechanical loading and hygrothermal exposure to induce creep, strain was observed to exhibit stepwise ‘‘jumps’’ (see Fig. 4) (34). The phenomenon of microcrack nucleation and propagation, which can increase the rate of degradation of composites, was identified both in cycling tests and in constant immersion tests (34–36). In agreement with the above discussion, changes in mechanism were observed to affect Fig. 3 Graph of the percentage weight of carbon/polystyrylpyridine samples following hygrothermal aging at 150°C of 250°C after drying (M0) and after 15 days of absorption (M m) as a function of the number of cycles. (From Ref. 33.) Long-Term Degradation of FRPs 429 Fig. 4 Creep of SMC-R57 glass/epoxy sheet molding compound at 90°C and 50% relative humidity under different loads (percent of static ultimate tensile strength). (From Ref. 34.) not only the degradation rate as a function of time at a given temperature but also the temperature dependence of the degradation process, causing deviations from simple Arrhenius-type behavior (37). Such limitations and complexities must be addressed in order to establish a basis for reliable prediction of the extent of environmental degradation of FRPs in long-term service. Further discussion of specific cases of supralinear behavior is given in the following sections. IV. EVALUATION OF ENVIRONMENTAL DEGRADATION: EXPERIMENTAL TECHNIQUES A. Mechanical Properties As the mechanical properties constitute the primary criteria for FRP performance, measurements of such properties have been widely used to characterize the extent of FRP degradation upon exposure to various environments. Practically every mechanical property has been used in such characterization, including tensile strength and modulus (38), compressive strength and modulus (20), flexural strength and modulus (28), Poisson’s ratio (38), short-beam shear strength (39), impact strength (40), and fracture toughness (41). Moduli are usually much less sensitive to environmental exposure than strength values because moduli reflect the structure of the entire specimen, whereas strength is determined by the region of the specimen most affected by exposure to environmental attack (32,42). However, moduli, too, can exhibit delayed increases in degradation rate after the specimen has become fully saturated with respect to moisture in the cases of plastics reinforced with fibers that are particularly prone to absorption of moisture (i.e., 430 Barkatt Fig. 5 Flexural modulus of Kev 49/T–300/Kev 49, an aramid-graphite/epoxy hybrid composite, as a function of moisture exposure (days) and temperature. (From Ref. 38.) aramid fibers) (see Fig. 5) (38). It has been recognized that the conditions under which samples are held between the end of the exposure period and the mechanical test have significant effects on the test results (43). During the intervening period, samples often undergo partial drying that causes further changes (decrease or increase) in the extent of degradation of strength resulting from the preceding exposure to moisture (43). This is one example of the difficulties in comparing results obtained under different test procedures and of the need to standardize such procedures. The procedures for testing of composites developed by the American Society for Testing and Materials (44) provide a useful starting point in controlling the test parameters. Although the results of mechanical tests provide data that are closely related to the properties that determine the performance of composites in service, such results are, in general, insufficient to provide a full basis for predictive modeling. As detailed earlier, the degradation kinetics has been found in many cases to be complicated and to involve the possibility of supralinear behavior at some stage of the exposure. This complexity requires a more thorough understanding of the microscopic processes of environmental degradation. Furthermore, such understanding is required in order to overcome other limitations of most mechanical tests. One such limitation is the difficulty in extrapolating results obtained for test coupons to full-size structural components. Another limitation involves tak- Long-Term Degradation of FRPs 431 ing into consideration the presence of other materials used along with FRPs. The presence of cement, for instance, is known to have significant effects on the degradation of FRPs based on glass fibers, and the presence of metals affects the degradation of FRPs based on carbon fibers. Due to all of these reasons, it has been found necessary, in many studies, to supplement the results of mechanical tests on FRP samples exposed to environmental degradation by measurements of various physical and chemical characteristics indicative of changes of the structure of the composite at the microscopic level. B. Thermal and Thermogravimetric Methods Gravimetric analysis of changes in sample weight following exposure to aqueous or humid environments has been widely used in monitoring the consequences of the exposure. In general, such exposure leads to a gain in weight, reflecting the uptake of water by the composite. Exceptions are observed in high-pH environments, where fiber dissolution offsets water absorption by the polymeric matrix (45). However, even in the absence of extensive fiber dissolution, interpretation of weight-gain data is not always straightforward. Ideally, such weight gain is expected to be linear and reversible. If the capacity of the matrix for uptake of water is constant, then the weight gain observed upon exposing the FRP for a long period of time should be linearly dependent on the relative humidity. Furthermore, upon bringing the sample back to an environment, which has the temperature and humidity at which the sample was held before the exposure, it is expected to return gradually to its original weight. In fact, sharp increases in moisture content have been observed upon exposure at high relative humidities (46,47) and the process of moisture uptake is not entirely reversible. These effects have been attributed to the formation of microcracks in the polymeric matrix (31). An even more serious limitation on the applicability of weight-gain measurements is the observation, based on 2-year immersion tests, that there is no direct correlation between the quantity of water absorbed and the loss of mechanical properties. Hence, absorption curves were concluded to be of limited use to material selection (48). Thermogravimetric analysis (TGA) has proven to be a very useful technique in evaluations of the effects of environmental exposure. TGA measurements at a heating rate of 10°C/min on a variety of polymeric composites previously exposed to various aqueous media showed that the weight loss observed upon heating from room temperature to 150°C reflected the expulsion of absorbed water. The weight loss observed between 150°C and 300°C was the most useful, reflecting the amount of monomer volatilization and thus of polymer degradation during the exposure. This weight loss also gave a good correlation with the extent of degradation of the mechanical properties such as the tensile, flexural, and shortbeam strength (42,45). TGA analysis can be supplemented by evolved gas analy- 432 Barkatt sis (EGA). Such analysis has confirmed that enhanced weight loss from FRPs previously exposed to aqueous environments is due to hydrolytic depolymerization (32). Dynamic mechanical analysis (DMA) and dynamic mechanical thermal analysis (DMTA) have been used to monitor damping properties, including the changes of these properties that occur at the glass transition (T g). These methods have been applied mostly to the polymeric matrix (49) but also to polymeric (aramid) fibers (50). As in the case of mechanical measurements, techniques have been developed to perform DMA measurements during immersion in order to avoid uncertainties due to changes upon sample storage following the exposure (51). The results of DMA and DMTA have been used to establish correlations between changes in T g and in mechanical properties, respectively (24,31,37,49,52). Very clear and quantitative measurements of changes in T g as a result of environmental exposure have been provided by differential scanning calorimetry (DSC). Determinations of T g by means of DSC have shown that exposure to water at elevated temperatures causes a decrease in T g due to partial hydrolysis and plasticization (42,45,53) of the polymeric matrix, which are also the factors involved in the enhanced TGA loss (see above). However, in the cases of commercial FRP matrices that are usually not fully cured, the decrease in T g due to depolymerization has been observed to be offset, in part or in full, by an increase due to additional cross linking (42). DSC measurements have also been used to measure changes in free volume upon thermal exposure (31). C. Microstructural, Spectroscopic, and Electrochemical Methods Optical spectroscopy has been extensively used to characterize major damage caused to FRPs by environmental exposure, including fiber buckling and matrix cracking (37,54). For a more sensitive characterization of microstructural damage, scanning electron microscopy (SEM) has been widely used to study failure modes and changes in failure mechanisms (33,54). Conventional SEM techniques have the disadvantage of requiring exposure of the samples to high vacuum, causing rapid and extreme drying that can enhance the actual damage by promoting shrinkage, cracking, and fragmentation. These artifacts can be minimized through the use of environmental SEM (E-SEM), which permits SEM measurements to be made in the presence of humid air at pressures only slightly below ambient. E-SEM measurements of polymeric composites have directly shown that the extent of the damage increases in aqueous media with high pH and at elevated temperatures (29), in agreement with the results of mechanical and thermochemical measurements. Long-Term Degradation of FRPs 433 X-ray diffraction (XRD) was used in microstructural studies of fatigue in glass fibers and their composites (55) and in studying the nature of composites formed on carbon-fiber composites during corrosion (13). Computerized image analysis was used to study surface degradation in glass fibers (56). Spectroscopic measurements provide particularly useful information about chemical changes in the molecular structure of the polymeric matrix as a result of exposure to corroding environments. Infrared (IR) spectroscopy has shown that epoxy matrices are susceptible to hydrolysis of phthalate ester linkages upon exposure to aqueous environments (14). The degradation of urethane linkages in polyurethane foam matrices was also followed using IR spectroscopy (57). Raman spectroscopy provides an alternative technique for monitoring changes in chemical bonding as a result of environmental exposure. Compared with IR spectroscopy, Raman scattering has the advantage of being less sensitive to the presence of absorbed water (32). Both IR and Raman measurements can be carried out using microprobes, permitting the use of these techniques to monitor chemical changes in particular areas of the composite, especially in the critical region constituting the fiber–matrix interphase. The use of such Raman techniques has given evidence of selective degradation of acrylic ester linkages in a vinylester–glass composite with observable release of acrylic acid monomers but no observable effect on the aromatic functional groups of the polymeric matrix (29). In addition to IR and Raman spectroscopy, electron spin resonance (ESR) spectroscopy was used to monitor the formation of free radicals upon bond cleavage in nylon fibers exposed to environments containing nitrogen oxides (58). Electrochemical potential measurements were used to follow the degradation of composites reinforced with carbon fibers exposed to electrolytic solutions, such as seawater, in the presence of metals (13,14). In such cases, the electrochemical potential gives a measure of the extent of galvanic interaction, which is responsible for the degradation of the carbon fibers. V. DEGRADATION OF FIBER MATERIALS A. Glass Because of their low cost, silicate glass fibers have the greatest potential of serving as reinforcements in FRPs to be used in large-scale civilian applications. The structure of such glasses is based on a three-dimensional network of silicate tetrahedra, modified through the introduction of alkali oxides (e.g., Na2O) and alkaline earth oxides (e.g., CaO) to reduce melt viscosity and permit processing at moderate temperatures (typically around 1200–1300°C). Multivalent oxides (e.g., Al2O3, B2O3) are also added in many cases in order to improve chemical resistance, and they may also improve the mechanical, thermal, or optical proper- 434 Barkatt ties. Thus, E-glass, the most common glass used in fiber reinforcements, consists of 54.3% SiO2, 17.2% CaO, 15.2% Al2O3, 8.0% B2O3, 4.7% MgO, and 0.6% Na2O by weight. The interaction between silicate glasses and aqueous environments has been extensively investigated since the beginning of the 20th century. It has been established that glasses have a very high chemical resistance to oxidants, reductants, and organic solvents (except amines, which have a basic character). Silicate glasses, unless they have a sizable content of oxides of heavy and multivalent metals such as Pb, Fe, and Al, have excellent resistance toward attack by acids (except HF). Commercial silicate glasses are very durable in aqueous media with near-neutral pH. The only environments in which glasses are subject to rapid attack are basic media (i.e., solutions with pH levels in excess of approximately 9) and hydrofluoric acid solutions. Recently, it has been shown that concentrated aqueous solutions of salts of alkali metals or calcium also attack glasses at moderately high rates (59). The cause of the high reactivity of alkali toward glass is the ability of the hydroxide ion, OH⫺, to break the siloxy bonds that hold together the glass structure. (In the case of HF, the fluoride ion fulfills a similar role.) If the starting medium is water, OH⫺ ions can be generated as a result of selective leaching of alkali or alkaline earth ions from the glass. This mechanism has been formulated in the following scheme (60): DSiEOENa ⫹ H2O → DSiEOEH ⫹ Na⫹ ⫹ OH⫺ (3) DSiEOESiD ⫹ OH⫺ → DSiEO⫺ ⫹ DSiEOEH (4) DSiEO⫺ ⫹ H2O → DSiEOEH ⫹ OH⫺ (5) The attack on Si sites in the surface regions of the glass continues until all four siloxy bonds are hydrolyzed and the resulting Si(OH)4 monomer passes into the solution (59,61). Remarkably, upon considering the interaction of water with glass, the simple mechanism consisting of reactions (1)–(3) can give rise to various types of kinetic behavior. As an illustration, several cases of glass–water interaction may be considered: 1. The glass is exposed to attack by water containing low levels of solutes with the volume of the water being sufficiently large and/or the water flowing at a sufficiently high rate to prevent significant accumulation of glass-dissolution products in the water. Leaching of alkali and alkaline earth species according to reaction (1) first proceeds rapidly, resulting in the buildup of a high-silica surface layer, which slows down further interaction of the water with the glass. Eventually, the depth of this layer reaches steady state and further attack on the glass proceeds slowly at a constant rate determined by the hydrolysis of the Long-Term Degradation of FRPs 2. 3. 4. 5. 435 siloxy network. The overall kinetics of glass corrosion is initially sublinear and later becomes linear with time (18). The volume of the aqueous environment in contact with the glass is limited, flow is very slow or nonexistent, and the solution is buffered against significant changes in pH. Under these conditions, dissolved silica levels eventually approach the saturation limit, resulting in almost complete retardation of further attack (62). The sublinear and linear stage of Case 1 are followed by another sublinear stage. In the cases of complex glasses, the saturation effect described in Case 2 may be followed by the formation of highly insoluble crystalline silicates. This results in lowering of the concentration of dissolved silica in the solution, elimination of the saturation effect, and renewed increase in the rate of glass dissolution (63). The volume of water is limited and flow is very slow, as in Case 2, but the aqueous medium is not buffered. As a result, the concentration of OH⫺ ions increases as reaction (1) proceeds OH⫺ ions consumed during attack on the siloxy bond network in reaction (2) are regenerated in reaction (3). Progressively higher pH results in increasing rate of attack on the glass. The sublinear and linear stages of Case 3 are followed by a supralinear stage (64). The glass is exposed to water, resulting in the formation of dealkalized, hydrated surface layer. Hydration is enhanced by factors that may be related to the glass composition [low silica content (65)], to the composition of the solution [high concentrations of alkali ions or Ca2⫹ (59)], or to environmental factors [high temperature]. This result in buildup of large hydration stresses. Eventually, the surface layer cracks and spalls off, destroying the retarding effect of this layer and opening up a large area of fresh glass for the surrounding water to attack. This results in a sharp increase in the effective rate of glass corrosion. This effect may be further promoted by certain exposure conditions such as wet–dry cycling. When a composite reinforced with silicate glass fibers undergoes penetration of moisture, the fibers become exposed to a very small volume of water which migrates very slowly into and out of the contact region. Wet–dry cycling is also likely to occur in many service environments. Accordingly, the most likely scenarios with respect to the fate of the fibers are Cases 2, 4, and 5. This implies that both sublinear/linear and supralinear kinetics of glass dissolution are possible. Both types of kinetics have indeed been observed. In particular, measurements of silica dissolution from fibers incorporated into FRPs have given evidence of very large increases in the effective dissolution rate with time (29,32). These increases cannot be solely attributed to the growing water content of the FRP (29). 436 Barkatt B. Carbon Carbon fibers are more chemically resistant than silicate glass fibers. In general, their solubilities and dissolution rates in both aqueous media and organic solvents are very low. In addition, they possess excellent mechanical and thermal properties. The main obstacle for the use of composites reinforced by carbon fibers in large-scale construction applications is economic. It should be noted, however, that despite their high resistance to chemical dissolution, composite materials reinforced with carbon fibers are susceptible to galvanic corrosion and other modes of electrochemical degradation when they are in contact with metals in electrolyte solutions such as seawater (13,14,66). C. Aramid Aramid fibers are known to be more susceptible in aqueous environments because, unlike silicate glass and carbon fibers, they absorb appreciable amounts of water. This can result in significant effects on their mechanical properties (36,39,67), causing a decrease in strength by as much as 40–70% depending on the temperature of the exposure (36). The degradation of the mechanical properties of polymeric fibers such as aramid fibers can be attributed to hydrolytic processes similar to these that take place with polymeric matrices (see Sec. VI). The effect of such processes on polymeric fibers is particularly noticeable because of the primary role that the fibers have in bearing the mechanical loads applied to the composite materials under service conditions. VI. DEGRADATION OF MATRIX MATERIALS Polymer degradation in various chemical environments has been extensively studied. The major effects of polar solvents, such as water and aqueous solutions (as well as alcohols, ammonia, hydrazine, etc.), on polymers can be described in terms of solvolytic reactions which cause the breaking of CEO or CEN bonds. For instance, polyester is held together by linkages consisting of carboxylic acid ester. This process can be described as (68) O 储 ECECEOECE → ECEOH ⫹ HOECE (6) Similarly, polyurethanes undergo solvolysis through the breaking of CEN bonds. Such solvolysis reactions are known to be catalyzed by acids as well as by bases: Long-Term Degradation of FRPs 437 O O 储 储 RCEOR′ → RC ⫹ HOR′ ⫹ H⫹ | ↑ ↑ OH HEOH H ⫹ (7) O O 储 储 RCEOR′ → RC ⫹ HOR′ ⫹ OH⫺ | ↑ ↑ OH⫺HEOH OH (8) Accordingly, accumulation of carboxylic acids formed upon hydrolysis of ester linkages in polymers can accelerate further hydrolysis, making it possible for the degradation kinetics to exhibit supralinear degradation kinetics. It has also been noted that in bulk polymers, the attack is initially restricted to the exposed surfaces and the rate of hydrolysis is related to the ability of the polymer to absorb water (68). A polymer serving as an FRP matrix and being interspersed with fibers is more susceptible to the cleavage of CEO or CEN linkages than the same polymer in bulk form without the presence of fibers. One reason for this is local buildup of OH⫺ ions (due to leaching of glass fibers) and H3O⫹ ions (due to the release of acidic groups upon hydrolysis of ester bonds). These ions accumulate within a confined volume adjacent to the polymer structure, where they can catalyze further hydrolysis. Another reason why hydrolysis of polymeric FRP matrices is expected to be more rapid than hydrolytic attack on bulk polymers is the fact that the fiber–matrix interfaces provide pathways for faster moisture absorption. Most FRPs are based on thermosetting polymeric matrices, including epoxy resins as well as polyester, vinylester, and less common polymers such as polystyrylpyridine (PSP). As mentioned earlier, exposure to moisture usually results in degradation of mechanical properties as well as in the value of T g (the glass transition temperature) (36,37,69), as a result of hydrolytic depolymerization. However, exposure to moisture at elevated temperatures can result, as mentioned earlier, in an increase in T g (70) due to enhanced cross linking in the cases of incompletely cured polymers (29). The change in T g may occur abruptly at the end of a long period over which no significant changes are observed (see Fig. 6) (71). Thermoplastic polymers, which are less commonly used as FRP matrices, are also affected by aqueous environments. Glass–polyphenylene sulfide (PPS) composites continue to have good mechanical properties upon thermal exposure even above T g, but when elevated temperature is combined with the presence of water, these properties deteriorate, due to degradation at the fiber–matrix interfaces (72). 438 Barkatt Fig. 6 Relative change in the glass transition temperature of AS/2220-3 carbon/epoxy composite as a function of exposure time to 1 atm and three different temperatures: 20°C (䊊), 90°C (䉭), and 140°C (䊐). (From Ref. 71.) Fig. 7 Master relaxation modulus curves for [⫾45] s T300/934 and GY70/339 carbon/ epoxy laminates. Horizontally shifted T300/934 data shown. (From Ref. 73.) Long-Term Degradation of FRPs 439 It is interesting to note that polymeric matrices exhibit viscoelastic behavior. At near-ambient temperatures, polymeric composites exhibit ‘‘elastic’’ or glassy behavior, but upon exposure to moisture at elevated temperatures, they experience transition from glassy to leathery behavior with a corresponding change in the time dependence of the modulus (see Fig. 7) (73). VII. INTERPHASE DEGRADATION The fiber–matrix interphase is a critical region in controlling the properties of FRPs and the degradation of these properties. Adhesion between the fibers and the matrix is widely recognized to be the result of a combination of mechanical effects and chemical bonding. On a microscopic scale, both the fibers and the matrix exhibit a certain degree of roughness. As a result, purely mechanical interlocking between the surface features on the fibers and on the matrix may provide a significant contribution to fiber–matrix adhesion. However, in general, such mechanical effects alone do not provide sufficient adhesion (74). A second major factor is wetting of the fibers by the matrix. Wetting requires minimization of the presence of impurities on the surface of the fibers and of trapped air or gas bubbles at the interface. The introduction of silane coupling agents generally improves fiber–matrix wetting. Finally, actual chemical bonding between the fibers and the matrix provides an important though variable contribution to the adhesion in composites. Such bonding is achieved through the use of organosilane coupling agents. The reaction of the coupling agent with the fiber and with the matrix can give rise to bonds of varying strength, ranging from strong covalent forces to weak van der Waals forces. Typically, the organic functional groups of such silanes provides covalent bonding with the polymeric resin, whereas the partially hydrolyzed silane group forms hydrogen bonds with the fibers (75). Exposure of composites to moderately elevated temperatures under dry conditions does not necessarily have a significantly detrimental effect on the three types of interaction involved in the adhesion between the fiber and the matrix mentioned. However, exposure to aqueous or humid environments eventually has to lead to a reduction in the strength of the adhesion. The bonds formed by silane groups, including hydrogen-type SiEO---HEOESi bonds and even covalent siloxane (DSiEOESiD) bonds, are highly susceptible to hydrolysis into separate silanol (DSiEOEH) groups. Eventually, corrosion of fibers may result in smoothing of sharp protuberances and jagged edges and thus reduce the extent of mechanical interlocking as well. Under mechanical loading, hydrolytic attack reduces the resistance of composites to debonding between the fibers and the matrix. This facilitates fiber pullout and cracking of the unprotected sections of the fibers, resulting in earlier failure than in the case of composites which have not been weakened by the effects of exposure to moisture. 440 Barkatt Various characterization methods, including mechanical, microstructural, and chemical methods, have been applied to the study of the interphase regions in FRPs (76). It has been shown that the absorption of moisture at the interphase affects fiber–matrix adhesion (69). The importance of treatments of the surface of the fibers on the performance of FRPs under service conditions has been shown (67). The correlation between the specific properties of the interphase region and the overall mechanical performance of FRPs has also been established (49). Enhanced corrosion of both the fiber and matrix materials in the interphase region has been demonstrated and the relevant mechanisms investigated (29). VIII. EFFECTS OF CORROSIVE ENVIRONMENTS RELEVANT TO CIVIL APPLICATIONS: GENERAL TRENDS A. Effects of Chemical Environment The effects of exposure to various environments that may be encountered in service on FRP materials have been extensively surveyed. Much of the work carried out in the 1970s and 1980s was summarized in three volumes edited by Springer (1–3). An important early study was carried out on the behavior of glass-reinforced polyester and vinylesters at two temperatures (23°C and 93°C) in a variety of environments, including humid air (50% and 100% relative humidity), saturated solutions of NaCl in water, diesel fuel, lubricating oil, antifreeze (a mixture made up of equal volumes of ethylene glycol and of water), and gasoline (16). Measurements over periods of up to 6 months showed that at 23°C, the effect of the aqueous NaCl solution on the tensile strength was largest. Significant reduction in tensile strength was also observed in 100% humid air and in the water– ethylene glycol mixture, whereas the effects of the nonaqueous solvents were generally not significant. The effects of the aqueous environments were considerably enhanced at 93°C, whereas the effects of the nonaqueous environments remained largely insignificant. Shear strength was affected to a smaller extent than tensile strength, but it too exhibited a large decrease upon exposure to an aqueous salt solution or to water–ethylene glycol mixtures. Moduli were affected much less than the corresponding strengths, reflecting the fact that overall strength is determined by the strength of the most degraded section of the specimen, whereas modulus is a manifestation of the average state of the entire specimen (see above). Measurements of weight change indicated widespread deviations from Fickian behavior. It was very interesting to note that in the case of 100% humid air, for instance, the kinetics of strength degradation appeared to be supralinear. No significant change in tensile strength took place at both temperatures at periods of up to 2 months, but a considerable drop was observed in tensile strength after Long-Term Degradation of FRPs 441 6 months, despite the fact that the absorption of moisture (as reflected in weight gain) was completed in 1 month at 23°C and in a few days at 93°C. In aqueous environments, the pH and the presence of solutes have a significant effect on fiber corrosion and, as a result, on the kinetics of degradation of the composite as a whole. A recent extensive study on vinylester–glass and polyester–glass composites (42,77) has shown that at high-pH fiber dissolution is much faster, as expected on the basis of Eqs. (2) and (3). This agrees with the well-known incompatibility of glass-based composite components with cement environments (27,78) which are characterized by highly basic pore water (15). Strong acid solutions also cause degradation of the strength of glass fibers due to ion exchange of protons for metal ions on the glass surface (79). As mentioned earlier, strong acids and strong bases also catalyze depolymerization of the polymeric matrix and hydrolysis of siloxy bonds in the interphase region. Aqueous solutions of strong acids, such as sulfuric acid, were shown to cause accelerated degradation as a result of interface corrosion (35) and formation of fractures (80). In the cases of aramid fibers, the fibers as well as the matrix undergo accelerated depolymerization in strongly acidic or strongly alkaline solutions. Weakly acidic solutions, such as acetic acid solutions buffered at pH values of 3–5, do not cause significantly more extensive corrosion than deionized water (77). Carbon fibers are the most resistant to acidic or alkaline environments, but higher concentrations of electrolytes accelerate galvanic corrosion in the presence of metals (see above). B. Effects of Temperature and Moisture Content Exposure of polymeric composites to temperatures exceeding 300°C results in massive depolymerization and volatilization (42,81) together with rapid degradation of the mechanical properties (24,82). Such temperatures, however, are encountered only in the vicinity of fire and are outside the range of normal service environments. At lower temperatures, simple Arrhenius-type behavior is observed only over limited temperature ranges. In particular, the degradation kinetics of FRPs can be generally expected to change once the temperature significantly exceeds the glass transition range of the polymeric matrix. In the case of epoxy–aramid composites, for instance, it was observed that even under dry conditions, mechanical properties such as flexural strength and stiffness were complicated functions of temperature. In the presence of moisture, the deviations from simple Arrhenius-type behavior are even more noticeable, with the strength exhibiting a sharp drop with increasing temperature below room temperature, then only a small dependence between room temperature and T g, and then a sharp drop above T g (83). An important factor involved in the complex degradation patterns was the observed strong dependence of T g itself on the moisture content 442 Barkatt of the composite, amounting to a drop of 60°C observed when the moisture content increased from 0% to 2%. These changes in T g are themselves a function of the time of exposure to elevated temperatures and moisture, and the time dependence often exhibits a supralinear dependence, with rapid variation following a period of little change (71). Because of the importance of changes in transport mode which take place at T g, the extent of degradation of mechanical properties is strongly dependent on the fraction of the material exhibiting the glass transition (i.e. on the degree of noncrystallinity) (69). In tests on laminates, large differences in the sensitivity with respect to exposure to elevated temperatures and to moisture were observed among laminates of different orientations among the fiber directions in adjacent laminas (84). In this study, it was noted that large scatter (in many cases, 20–60%) existed in the data obtained for the degradation of the mechanical properties. In addition, it was noted that the loading rate in the mechanical tests and that interactions among this rate, the temperature, and the moisture content were also expected to be significant. The importance of considering the ply and the laminate as well as individual fiber–matrix elements in degradation studies has been stressed (24). Even more complex behavior was observed when FRPs were exposed to periodically varying environmental conditions. Thermal cycling at relatively low temperatures was shown to result in cracking of the polymeric matrix (85), especially when it involved freeze–thaw cycles. Thermal cycling between long exposures at 70°C in humid air and short exposures at 150°C was shown to give rise to very prominent supralinear behavior, with the mechanical properties starting to change rapidly after approximately 1 year (33). This was attributed to microcracking due to the alternation between a high moisture content at the lower temperature and a low moisture content at the higher temperature. C. Gaseous and Radiation Environments Exposure of aramid–epoxy composites to gaseous environments containing air pollutants such as nitrogen oxides was shown to result in degradation of mechanical strength as a result of the formation of free radicals capable of attacking the polymeric fibers (86). At elevated temperatures, the diffusion of atmospheric oxygen and its attack on the matrix become important (24). In addition to chemical environment and temperature, short-wavelength radiation is also known to affect the properties of polymeric materials. Degradation of marine fabrics containing polyester or nylon fibers was observed upon exposure to ultraviolet radiation under dry or wet conditions (87). The mechanism of degradation is likely to involve the formation of free radicals, similar to the case of attack by nitrogen oxides. Bond scission and the formation and propagation of free radicals also constitute the mechanism of thermal degradation and depolymerization of polymeric materials in general under dry conditions (68). Long-Term Degradation of FRPs 443 D. Stress Corrosion A number of studies were carried out on the degradation of FRP materials under a combination of environmental exposure and mechanical loads (35,37,77,80,86). In general, moderate mechanical loads were observed to cause a moderate increase in the severity of the effects of environmental attack. However, definite dependence of the rate of crack growth on stress intensity was observed. It was concluded that the extent of stress corrosion reflects the susceptibility of the fibers to corrosion and to shrinkage in the corroding environment. Stress corrosion was observed to depend on the temperature, the acidity of the medium, and the type of glass fibers used as reinforcements (88). The dependence of crack growth velocity in glass on stress intensity and the effects of the environment on this dependence have been thoroughly investigated (89,90). The velocity of crack propagation was found to be controlled in its initial and most important stage by the rate of stress corrosion at the crack tip, according to the equation V⫽ Ax0n exp(bK I) n (9) where V is the velocity of crack propagation, x0 is the partial pressure of water (i.e., the relative humidity), n is the order of the chemical reaction (varying from 0.5 to 1 with increasing relative humidity), K I is the stress intensity factor, and A and b are constants. The initial stage (region I) is followed by a stage (region II) where the crack-propagation velocity is controlled by the rate of diffusion of moisture to the crack tip, independent of stress intensity: V⫽ CD(H2O)x0 δn (10) where D(H2O) is the diffusivity of water in the environment, δ0 is the boundarylayer thickness, and C is a constant. In the third and final stage (region III), the crack-propagation velocity again exhibits a strong dependence on the stress intensity, but is no longer dependent on the relative humidity. This mechanism has to be taken into account when glass fibers are used as FRP reinforcements. In the cases of FRP materials, delamination crack propagation has been shown to exhibit classical supralinear behavior leading to failure upon exposure to elevated temperatures under conditions of either static creep (see Fig. 8) or cyclic creep (see Fig. 9) (91). Measurements of crack growth in sheet-molding compounds have shown that the dramatic increase in crack growth occurs once a critical crack length is achieved. The critical crack length (see Fig. 10) is a constant characteristic of each sheet-molding compound. Both the time to achieve the critical crack length and the subsequent rate of crack growth are functions of the applied stress intensity, K I (92). As noted earlier, under conditions conduc- 444 Barkatt Fig. 8 Crack-propagation curve of a carbon/PEEK composite in static creep at 180°C. (From Ref. 91.) Fig. 9 Crack-propagation curve of a carbon/PEEK composite in cyclic creep at 200°C. (From Ref. 91.) Long-Term Degradation of FRPs 445 Fig. 10 Crack growth in a glass/polyester sheet molding compound as measured by direct visual observations. (From Ref. 92.) tive to creep, stepwise ‘‘jumps’’ in strain may occur at random times (34). The sharp increase in creep observed in bonded FRP joints (see Fig. 11) was found to occur somewhat sooner at higher loads (see Fig. 12) (93). E. Effects of Hydrostatic Pressure The corrosion of silicate glass is not significantly affected by hydrostatic pressure (94). However, in the case of polymeric composites, pressure can affect the free volume and, thus, the effective T g (33). The saturation moisture content of the graphite–polymer composite was shown to increase under large hydrostatic pressures although the diffusion coefficients remained unchanged (28). Moisture absorption in epoxy composites is accelerated under large hydrostatic pressures, although the effect of pressure varies greatly from one composite to another (48), possibly due to variations in polymeric structure or microporosity. Thus, the main effect of large hydrostatic pressures appears to consist of squeezing more moisture into the polymeric matrix and, possibly, the interphase. 446 Barkatt Fig. 11 Creep deformation in a glass/polyester sheet molding compound as a function of time. Load levels indicated are percent of baseline value. Solid circles indicate test coupon failed. (From Ref. 93.) Fig. 12 Creep deformation in a glass/polyester sheet molding compound as a function of time. Load levels indicated are percent of baseline value. (From Ref. 93.) Long-Term Degradation of FRPs 447 IX. MECHANISMS AND MODELS In many studies, it has been attempted to develop quantitative models for the degradation of the mechanical properties of FRP materials as a result of exposure to various environments, in particular under hydrothermal conditions combining moisture and moderately elevated temperatures. Many of these models are based on Fickian diffusion of moisture into the composite as the rate-controlling process. However, as mentioned earlier, the uptake of moisture is, in many cases, non-Fickian (95). The saturation level with respect to moisture may change with time, and degradation in mechanical properties often exhibits large changes in rate after the moisture content has leveled off (16). The detailed kinetics of degradation is therefore highly dependent on the specific material under investigation and on the nature of the corroding environment. Accordingly, the applicability of such mechanical models is limited, in general, to specific combinations of material and environment. Chemical models of FRP degradation are usually more phenomenological and less mathematical in nature. Such models attempt, for instance, to describe the interaction of FRP materials with their environment in terms of changes in structural parameters such as T g (71,96). Such chemical models take into consideration mechanistic features such as ion exchange on fiber surfaces (80), chemical interactions and physical absorption of water into the polymeric matrix (39), and hydrolytic depolymerization and thermally induced cross-linking (85,95). Such chemical models require good understanding of the detailed microscopic mechanisms involved in the interaction between the corroding environment, on one hand, and the fibers, the matrix, and the interphase, on the other. A. Fiber-Based Mechanisms Studies of fiber-based degradation mechanisms have shown a large variety in mechanisms among different systems and exposure conditions. Corrosion degradation of fibers has been observed in cement environments, which produce an alkaline aqueous phase that is highly corrosive toward glass and aramid fibers (see Section VIIIA). Once this solution penetrates through the matrix, rapid attack on the fibers begins, leading to eventual failure (97,98). Of course, this phenomenon is expected to result in a rise in corrosion rates at the time that the corrosive medium reaches the fibers. Acidic environments also give rise to failure mechanisms resulting from alteration of fiber surfaces due to exchange of protons for metal ions (80). Fiber-based failure mechanisms have been shown to change from direct tensile rupture at low temperatures and low moisture content to local fiber buckling and shear failure following matrix plasticization (36,37,99) or matrix cracking (54,100). 448 Barkatt B. Matrix-Based Mechanisms Matrix degradation usually does not lead to failure directly, but opens the way to environmental attack on the fibers (36,97–99) or on the interphase (67), resulting in ultimate failure. The principal mechanisms of matrix degradation consist of microcracking at relatively low temperatures (below T g) and low moisture content and of softening at elevated temperatures (above T g) and high moisture content (which tends to cause plasticization, resulting in a lower T g) (54). As detailed in the previous paragraph, these modes of matrix degradation are followed by the onset of fiber degradation and eventual failure. Of course, at the time when the rate-determining step changes from matrix degradation to fiber corrosion, the rate of overall degradation is expected to exhibit a very significant change. This greatly complicates extrapolation of short-term results to longer times. Furthermore, the change in failure mechanisms with increasing temperature also complicates the development of models intended to account for the temperature dependence of the degradation process. As noted earlier, the controlling failure mechanism at low temperatures and low moisture content is often matrix cracking and tensile fiber failure, whereas at elevated temperature and high moisture content, the dominant mechanism becomes matrix plasticization and softening opening the way to fiber buckling and failure in shear. This change makes it impossible to use a simple Arrhenius-type model to predict the degradation behavior over a broad temperature range or to evaluate the behavior of FRPs with high moisture content from measurements performed while the moisture content is low. C. Interface-Based Mechanisms As noted earlier, the interphase is a region that is particularly prone to moisture ingress and to degradation because of the relative weakness of the bonds in this region and their susceptibility to hydrolysis. The mechanisms of interphase degradation involve loss of adhesion due to such hydrolytic attack (33,35,69,72). This explains why pretreatment of the fibers has an important effect on the strength of FRPs (e.g., aramid–epoxy composites), both before and after exposure to moisture and to elevated temperatures. However, increasing moisture concentration eventually reduces the beneficial effects of such pretreatment (67). D. Multistage Mechanisms A recent detailed investigation of the chemical steps involved in the degradation of glass–vinylester composites (29) showed that ingress of water leads to leaching of alkali ingredients from the fibers and to a local rise in pH around the fibers, which was identified by means of the use of acid–base indicators. Elevated pH results in enhanced corrosion of the fibers and leads to degradation and perforation of the interphase, allowing further ingress of water. Hydrolytic attack on the Long-Term Degradation of FRPs 449 surrounding matrix was shown to result in depolymerization and loss of acrylic acid monomers, as manifested in micro-Raman spectroscopic measurements and in measurements of the extraction of acrylic acid into the surrounding water. Interface cracking and degradation of the matrix surrounding the fibers were directly observed using E-SEM. The effects of cracking and increased porosity were demonstrated by monitoring the concentration of dissolved silica in the aqueous phase and showing that the rate of dissolution is directly related to the magnitude of the exposed area, including cracks, open pores, and so forth. The overall result of this combination of processes is for the initial damage to open the way to more rapid subsequent degradation. This was reflected in a dramatic rise, by as much as one to two orders of magnitude, in the rate of dissolution of silica from the samples, which was observed after periods ranging between several weeks and several months at temperatures between 60°C and 80°C. At lower temperatures, such supralinear behavior may become evident after several years. This poses serious limitations on the prediction of long-term performance. X. CONSIDERATIONS OF SCALE AND MATERIALS VARIABILITY A. Scale Problems In the cases of homogeneous, dense materials such as metals and glasses, the size of structural components affects only the extent of corrosion which may be safely tolerated, but not the kinetics of the corrosion process itself. In the cases of FRP materials, on the other hand, size is an important consideration, because degradative processes involve internal surfaces as well as external surfaces. Obviously, saturation of a thicker slab or tendon of a composite with respect to moisture takes longer than the penetration of moisture through a thinner structural component. This would extend the period required for hydrolytic processes to affect the entire thickness of the material. It has been noted that laboratory coupons are not satisfactory for determining structural performance, as the load path is not always obvious because of the variation in the material properties due to the anisotropic nature of the composite (101). These considerations limit the range of long-term extrapolations and require coupon tests to be supplemented by tests on full-size structural components (102). Measurements on the transverse tensile strength of graphite–epoxy composites have shown that the strength decreased as the volume of the materials under stress increased (103). B. Materials Variability In the characterization of various properties of FRPs related to engineering applications, in particular resistance to environmental degradation, the issue of uniformity and standardization poses very serious problems. The fabrication of large- 450 Barkatt scale FRP products is generally performed under poorly controlled conditions. It was pointed out that ‘‘the quality manufacture of FRP, while quite well understood and practiced in quality FRP shops, is sometimes still on a steep learning curve in the construction field’’ (104). Many fabrication techniques for FRP products exist (105), and the chemical durability of the products differs widely from method to method and from manufacturer to manufacturer. The fabrication of typical polymeric matrix composites, such as glass–vinylester FRPs, involves complex materials systems combining glass fibers, resin systems including vinylester and styrene monomers, catalyst systems including catalysts as well as promoters and gel time retarders, fillers, and thixotropes such as fumed silicas, fire retardant additives, pigments, and mold-release agents (104,106). The relative amounts of these components and the exact fabrication procedure are subject to significant unintentional fluctuations as well as to controlled changes in response to variations in the fabrication conditions (temperature, humidity, equipment, age of ingredients). In contrast with the fabrication of FRPs for specialty applications (particularly in aerospace technology), the production of FRPs for civil engineering applications still lacks tight specifications and standards. One particular property of FRPs which has a very important effect on the resistance to environmental degradation is the porosity. Composites with significant porosity can be expected to be much more susceptible to moisture uptake and to subsequent hydrolytic attack even upon using more durable resins and fibers. Measurements on two commercial vinylester–glass composites following exposure to deionized water at 80°C (29) showed that the extent of acrylic acid extraction in the case of one of these materials was consistently higher by an order of magnitude than in the case of the other one. (The pH of the aqueous phase in contact with the first material was 3.4 ⫾ 0.2, and the pH of the aqueous phase in contact with the second material was 4.4 ⫾ 0.1.) Thus, questions regarding issues such as the relative chemical durabilities of FRPs based on vinylester, polyester, and epoxy cannot be answered unambiguously, because it is practically impossible to find FRP products made under identical conditions except for the identity of the resin. The differences in chemical durability among various vinylester-based FRPs, particularly those with wide variations in porosity, can be far greater than the differences between a particular vinylester composite and a polyester-based product. Furthermore, even the terms ‘‘vinylester,’’ ‘‘polyester,’’ and ‘‘epoxy’’ denote general classes of resins with widely varying compositions and molecular weights rather than specific chemical compounds. As detailed earlier, significant variations in composition, fabrication, and properties exist even among different lots of a single product produced by a single manufacturer. Accordingly, predicting the long-term characteristics of a certain FRP material using samples not taken from the actual components used in a particular project can introduce major uncertainties, whereas insistence on conducting lengthy tests with samples of the material actually in use entails very cumbersome and costly experimentation. Long-Term Degradation of FRPs 451 XI. PREDICTION OF LONG-TERM BEHAVIOR AS A LIMITING FACTOR IN EXPANDING FRP APPLICATIONS: CURRENT STATUS Uncertainties in predicting the long-term performance of FRP materials in largescale construction applications constitute a major impediment to rapid expansion of the use of FRPs in such applications, in spite of their attractive mechanical and chemical properties and their economic affordability. The methodology and database for predicting the long-term behavior of FRPs still lag far behind those available for other construction materials such as metals and cementitious materials. As detailed earlier, such predictive capability requires the existence of accepted methodology for carrying out relevant tests under service conditions, as well as accepted accelerating factors and accelerated test techniques for predictive testing. In addition, it is necessary to have reference materials that can be used to verify the validity of standard test procedures with a high degree of reproducibility and adequate accuracy. Another essential requirement is sufficient understanding of the environmental degradation mechanisms in order to make possible the use of short-term test results obtained under service conditions and under accelerated conditions in developing reliable predictions of long-term behavior. Substantial progress in these areas has been made in recent decades, but much more work has to be done toward meeting these requirements. In general, it has been established that simple kinetic laws, such as Fickian absorption isotherms of moisture absorption (107) and Arrhenius-type temperature dependence of hydrolytic degradation (27) may hold over limited ranges of time, temperature, and moisture content. Temperature has been established as the most promising accelerating factor, in preference to composition of the aqueous phase or mechanical loading. An increase in temperature from 23°C to 80°C, for instance, was found to leave unchanged both the relative effects of various aqueous environments on a particular composite, or the relative resistance of various FRPs toward a given chemical environment (108). However, these conclusions have been found applicable over a limited range of combinations of materials and environmental conditions, and over testing periods of up to several months. It was noted that the increase in degradation rate associated with a particular increase in temperature widely varies from one FRP material to another (48). Variations of temperature acceleration, such as raising the test temperature by a value corresponding to the lowering of T g due to the presence of moisture, have been proposed (20), but such methods are also limited to the temperature range over which the controlling mechanism remains unchanged. As detailed in the previous sections, extrapolations based on these simple models are bound to fail once the length of exposure, temperature, or moisture content produce a change in the nature of the rate-determining degradation process (20,37). Changes in mechanism that cause an increase in rate (supralinear behavior) are of particular concern. In the preceding sections, a large number of phenomena leading to 452 Barkatt changes in the controlling mechanism and, hence, in the degradation kinetics, after a certain period of exposure were discussed. Such phenomena include, for instance, the following: • A change in the chemical environment of glass fibers due to leaching of basic components, resulting in a local rise in pH and corrosivity • Cracking and spalling of the surface of glass fibers, causing an abrupt rise in the exposed surface area and the effective rate of corrosion • Extraction of acidic monomers produced upon hydrolytic depolymerization of the matrix, facilitating further matrix hydrolysis • Hydrolytically induced cracking in the matrix, increasing the void volume and thus the capacity for absorption of moisture • A change of the rate-controlling process from matrix degradation to fiber corrosion • A hydrolytically induced decrease in the T g of the polymeric matrix, facilitating migration of moisture toward the fibers • A change of the rate-controlling process from matrix degradation to interfacial debonding • Debonding and cracking in the interphase, opening up larger areas for hydrolytic attack These phenomena may cause changes (gradual or abrupt) in the observed degradation rate with time. In addition, other phenomena discussed earlier cause deviations from simple, Fickian dependence of the degradation process on the temperature and the moisture content. Such phenomena include the following: • An increase in the extent of cross-linking in the polymeric matrix at elevated temperatures • A change in the matrix degradation mode from cracking at a low temperature (below T g) and a low moisture content to plasticization and softening at a high temperature (above T g) and a high moisture content • A change in the overall failure mechanism from tensile rupture of the fibers at a low temperature and low moisture content to a shear failure resulting from fiber buckling following matrix plasticization at a high temperature and a high moisture content The frequent observation of changes in the nature of the rate-controlling mechanism and the potential for supralinear behavior in the degradation kinetics are major issues that need to be resolved in the quest for reliable prediction of the performance of FRPs in long-term service. The contribution of such phenomena to the extent of degradation has to be fully understood and bounded. Such understanding is also needed in order to identify accelerating factors for use in predictive testing and in specifying limits for the values of such accelerating factors to avoid encountering significant changes in mechanism. Long-Term Degradation of FRPs 453 Another extremely important area, in which the progress made to date has been inadequate, is the issue of representative sampling and testing. On one hand, FRPs used for large-scale applications at the present time are not well specified with regard to their composition and fabrication technique. On the other, there is almost no information as to the effect of moderate variations in composition and in fabrication technique on the resistance to environmental degradation. Accordingly, efforts to quantify such effects have to be undertaken, along with attempts to introduce tighter control and better specification of FRP products and to develop suitable reference materials. Although a promising start has been made in the development of standard test methods (44), such methods have to be further developed to make them applicable to large-scale structural components. For this purpose, availability of suitable reference materials is indispensable. Until these requirements are met, it is necessary to limit the use of predictive testing [e.g., to extrapolation by no more than one decade in time (102)]. Such limitations are likely to continue to slow down the introduction of FRPs into use in large-scale structural applications until the scientific and engineering issues related to longterm prediction are better resolved. ACKNOWLEDGMENTS The author is extremely grateful to Catherine R. Lang and Christina L. Monkres for very valuable help in preparation and organization of the manuscript. REFERENCES 1. GS Springer, ed. Environmental Effects on Composite Materials, Vol. 1. Westport, CT: Technomic Publishing Co., 1981. 2. GS Springer, ed. 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Mechanical alloying (i.e., solid-state mixing) is also effective for the preparation of amorphous alloy powders. Vitrification of metal surfaces is also made by the destruction of the long-range atomic order in the surfaces of solid metals. They are free of defects associated with the crystalline state, such as grain boundaries, dislocations, and stacking faults. Furthermore, the formation of the structure with no long-range atomic order is based on the prevention of solidstate diffusion during solidification to form equilibrium phases and, hence they are free of compositional fluctuations formed by solid-state diffusion, such as second phases, precipitates, and segregates. The amorphous alloys are, therefore, regarded as ideally chemically homogeneous alloys composed of thermodynamically metastable single-phase solid solutions supersaturated with alloy constituents. The formation of the single-phase solid solution supersaturated with alloying elements is quite suitable in producing new alloys possessing specific properties by alloying. Even if amorphous single-phase alloys are not formed alloys prepared by amorphization methods are often composed of nanocrystalline phases supersaturated with alloying elements. From a corrosion point of view, they can be considered as homogeneous alloys. 459 460 Hashimoto II. CORROSION-RESISTANT ALLOYS IN AQUEOUS SOLUTIONS The corrosion behavior of amorphous alloys has received particular attention since the extraordinarily high corrosion resistance of amorphous iron–chromium–metalloid alloys was reported in 1974 (1). Most of amorphous alloys prepared in the 1970s and in the first half of the 1980s were rapidly quenched from the liquid state (i.e., melt spinning). Accordingly, the corrosion behavior of meltspun amorphous alloys were extensively studied. The preparation of amorphous iron-based alloys by melt spinning generally requires the alloys to contain large amounts of metalloids which are mostly close to the eutectic compositions. The corrosion rate of amorphous iron–metalloid alloys decreases with the addition of almost all the second metallic elements such as titanium, zirconium, vanadium, niobium, tantalum, chromium, molybdenum, tungsten, cobalt, nickel, copper, ruthenium, rhodium, palladium, iridium, and platinum (1–9). A. High Corrosion Resistance of Amorphous Fe–Cr–Metalloid Alloys The addition of chromium is particularly effective in enhancing corrosion resistance. For instance, amorphous Fe–8 Cr–13 P–7 C alloy passivates spontaneously even in 2M HCl at ambient temperature (10). (The numbers denoting the concentrations of the alloy elements in amorphous alloy formulas are all expressed as atomic percent unless otherwise stated.) Amorphous Fe–3 Cr–13 P– 7C alloys containing 2at% molybdenum, tungsten, or other metallic elements are passivated by anodic polarization in 1M HCl at ambient temperature (11). The chromium addition is also effective in improving the corrosion resistance of amorphous cobalt–metalloid (12) and nickel–metalloid (13,14) alloys as shown in Fig. 1 (14). A combined addition of chromium and molybdenum is more effective. Some amorphous Fe–Cr–Mo–metalloid alloys passivate spontaneously even in 12M HCl at 60°C. Critical concentrations of chromium and molybdenum necessary for spontaneous passivation of amorphous Fe–Cr–Mo–13 P–7 C and Fe– Cr–Mo–18 C alloys in hydrochloric acids of various concentrations and temperatures are shown in Fig. 2 (15). In strong acids with high oxidizing power such as boiling nitric acids, the alloys with corrosion resistance based mostly on the presence of chromium are corroded, but amorphous alloys containing valve metals such as tantalum show very high corrosion resistance, which is much higher than that of crystalline tantalum metal as shown in Fig. 3 (16). Amorphous and Nanocrystalline Alloys 461 Fig. 1 Changes in corrosion rates of amorphous Fe-, Co-, and Ni-based metalloid alloys containing chromium in 1M HCl at 30°C as a function of alloy chromium content. (From Ref. 14.) B. Factors Determining the High Corrosion Resistance of Amorphous Alloys 1. Passive Films Rich in Cations of Alloying Elements with High Passivating Ability X-ray photoelectron spectroscopic study (17) of the spontaneously passive amorphous Fe–10 Cr–13 P–7 C alloy in 1M HCl has revealed that the passive film consists of Cr 3⫹, O 2⫺, OH⫺, and H 2O; hence, the passive film has been called a passive hydrated chromium oxyhydroxide film [CrOx (OH)3⫺2xnH2O]. Subsequent investigations have shown that the chromium enrichment occurs in passive films formed not only on amorphous alloys (13–15,18–20) but also on crystalline alloys (21–23) when their corrosion resistance is based on the presence of chromium. It has been known (24) that the resistance to passivity breakdown is higher when the chromium content of the passive film is higher. Accordingly, when an alloy has a higher ability to concentrate chromic ions in the passive film, the alloy has a higher corrosion resistance. The concentration of chromic ions in 462 Hashimoto Fig. 2 Critical concentrations of chromium and molybdenum necessary for spontaneous passivation of amorphous Fe–Cr–Mo–13 P–7C and Fe–Cr–Mo–18 C alloys in hydrochloric acids of various concentrations and temperatures. (From Ref. 15.) passive films formed on amorphous alloys is far higher than that on crystalline alloys, as shown in Table 1. The passive films formed by additions of sufficient amounts of valve metals to amorphous nickel–valve metal alloys are exclusively composed of valve metal oxyhydroxides such as TaO 2(OH) and NbO 2(OH) (16). Consequently, amorphous alloys containing strongly passivating elements, such as chromium and tantalum, have the very high ability to concentrate the beneficial ions in their passive films and have high corrosion resistance based on spontaneous passivation. 2. Homogeneous Nature of Amorphous Alloys The high corrosion resistance of amorphous alloys disappears using heat treatment for crystallization (25–30). Figure 4 shows an example of the effect of heat treatment (28). A nanocrystalline metastable phase is formed in the amorphous matrix by heat treatment at 703 K for 100 min. The alloy is no longer spontaneously passive in 1M HCl as soon as the nanocrystalline phase appears in the amorphous matrix, and the anodic dissolution current continues to increase with increasing time of heat treatment. This is due to the introduction of chemical heterogeneity into the amorphous alloy consisting of the chemically homogeneous single phase. Amorphous and Nanocrystalline Alloys 463 Fig. 3 Corrosion rates of amorphous Ni–Nb, Ni–Ta, and Ni–Nb–Ta alloys in boiling 9M HNO3 with 100 ppm Cr6⫹ ions. (From Ref. 16.) The detrimental effect of nanocrystalline heterogeneity in the amorphous matrix can be seen in Fig. 5 (29). Alloys were prepared by melt-spinning of liquid alloys composed of mixtures of eutectic Cr–13 P and Ni–19 P. The 59.81 Cr– 25.31 Ni–14.88 P, 58 Cr–27 Ni–15 P, 55.93 Cr–28.93 Ni–15.14 P, and 43.5 Cr–40.5 Ni–16 P alloys are identified as the amorphous structure by x-ray diffraction, but the former two alloys show more than three orders of magnitude higher corrosion rates than those of the latter two alloys. Detailed examination reveals that only a 0.14 at% decrease in the phosphorus content and only a 2.07 at% increase in the chromium content result in precipitation of the nanocrystalline 464 Hashimoto Table 1 Concentrations of Chromic Ion in Passive Films Formed on Amorphous Alloys and Stainless Steels in 1M HCl at Ambient Temperature Amorphous alloy Fe–10 Cr–13 P–7 C Fe–3 Cr–2 Mo–13 P–7 C Co–10 Cr–20 P Ni–10 Cr–20 P Stainless steel Fe–30 Cr–(2 Mo) Fe–19 Cr–(2 Mo) Cr3⫹ /total metallic ions Passivation Ref. 0.97 0.57 0.95 0.87 Spontaneous Anodic polarization Spontaneous Spontaneous 18 11 19 20 0.75 0.58 Anodic polarization Anodic polarization 22 23 body-centered-cubic (bcc) chromium phase of about 10 nm in diameter in the amorphous matrix and that preferential dissolution of the bcc precipitates takes place, although phosphorus-containing phases are not corroded. The formation of nanocrystalline phase in the amorphous matrix is not always detrimental. As will be mentioned later, amorphous Cr–Zr alloys consisting only of corrosion-resistant metals are extremely corrosion resistant in concentrated hydrochloric acids due to spontaneous passivation forming a double oxyhydroxide film of Cr3⫹ and Zr4⫹, and increasing chromium content increases the passivating ability and corrosion resistance. However, as an inherited characteristic from zirconium, the zirconium-rich alloys suffer pitting by anodic polarization. The change in the pitting susceptibility with structural change by heat treatment was examined (30). Specimens were heated with a rate of 4°C/min to the prescribed temperature, kept at the temperature for 30 min, and furnacecooled. As an example, polarization curves of Cr–67 Zr alloy specimens heated at different temperatures are shown in Fig. 6. These specimens are spontaneously passive in 6M HCl and their pitting potentials increase with heat-treatment temperature. The specimen heated to 500°C exhibits the highest pitting potential, but that heated to 600°C is lower than those of specimens heated to 400°C and 500°C. The heat treatment results in the formation of the less corrosion-resistant hexagonal closest packed (hcp) zirconium precipitates and the size of precipitates increases with heating temperature. The formation of the hcp zirconium phase leads to an increase in the chromium content of the matrix phase and hence to an enhancement of the formation of a thinner and more protective chromiumrich passive film covering the entire heterogeneous alloy surface. However, when the average size of the less corrosion-resistant zirconium phase exceeds a critical size (i.e., 20 nm), the protective chromium-rich passive film cannot completely cover the precipitates and the pitting resistance decreases. In this manner, if the precipitation of nanocrystalline phase enhances the passivating ability of the ma- Amorphous and Nanocrystalline Alloys 465 Fig. 4 Change in polarization curve of an amorphous Fe–10 Cr–13 P–7 C alloy in 1 M HCl with the time (min) of heat treatment at 703 K. (From Ref. 28.) trix, the precipitation of nanocrystalline phase is not always detrimental but sometimes increases the corrosion resistance. 3. High Activity of Amorphous Alloys When the chromium-enriched passive film is formed on amorphous and crystalline iron–chromium alloys, the composition of the alloy surface just under the chromium-enriched passive film is almost the same as that of the bulk alloy (22), although if nickel is contained in alloys such as austenitic stainless steels, nickel is concentrated in the underlying alloy surface because nickel is not contained in the passive film (31). Hence, the formation of the chromium-enriched passive 466 Hashimoto Fig. 5 Change in corrosion rate of melt-spun Cr–Ni–P alloys in 6M HCl at 30°C. The alloys were prepared changing the ratio of eutectic Cr–13 P to Ni–19 P. (From Ref. 29.) Fig. 6 Potentiodynamic polarization curves of sputter-deposited Cr–67 Zr alloy specimens measured in 6M HCl at 30°C. Specimens were heated at a rate of 4°C/min to the temperature indicated in the figure and kept at the temperatures for 30 min. (From Ref. 30.) Amorphous and Nanocrystalline Alloys 467 film results from selective dissolution of alloy constituents unnecessary for the passive film formation. When an alloy is able to passivate, fast active dissolution of the alloy results in rapid enrichment of beneficial ions. The passivating ability is, therefore, closely related to the activity of the alloy. Because amorphous alloys are thermodynamically metastable, unless passive films are formed, the amorphous alloys dissolve more rapidly than the crystalline counterparts. This can be seen from an example shown in Fig. 7 (32). The amorphous Co–25 at% B alloy does not contain elements forming a stable passive film in acids; hence, the dissolution rate of the alloy in an acid reflects the reactivity of the alloy. When heat treatment was conducted for 5 h, the amorphous Co– 25 at% B alloy begins to crystallize at temperatures higher than 523 K, forming a Co3 B phase in the amorphous matrix. The Co3B phase has the same composition as the amorphous matrix. The formation of the crystalline phase whose composition is the same as that of the amorphous matrix results in a decrease in anodic dissolution current. This reveals that the amorphous phase is more active than the crystalline counterpart. At 623 K, where the amorphous phase almost Fig. 7 Dissolution current density of Co–25 B alloy specimens measured in 0.5M Na 2SO 4 solution at pH 1.8 and 298 K where the alloy dissolves actively. Specimens were heat treated for 5 h at prescribed temperatures. (From Ref. 32.) 468 Hashimoto disappears, the single Co3B phase alloy shows the lowest activity for alloy dissolution. At even higher temperatures, the Co3B phase is decomposed to Co and Co2B phases by the disproportionation reaction. The formation of the heterogeneous three-phase structure leads to an increase in the alloy dissolution rate. The high reactivity of amorphous alloys based on the thermodynamically metastable nature is effective in enhancing the accumulation of beneficial passivating elements on the alloy surface as a result of the fast dissolution of unnecessary elements into environments and, hence, is responsible for fast passivation by the formation of the film in which the beneficial ions are highly concentrated, as shown in Table 1. The influence of the activity of alloys on the passivating ability can also be seen in Fig. 1. Among iron-, cobalt-, and nickel-based alloys, when the chromium content of alloys is not high enough to passivate spontaneously, the most active iron-based alloys dissolve most rapidly and the most noble nickel-based alloys dissolve most slowly. However, the fast active dissolution of iron-based alloys is effective in concentrating chromic ions on the surface and, accordingly, the iron-based alloys passivate spontaneously with the addition of the lowest amount of chromium. By contrast, the slowly dissolving most noble nickel-based alloys require the addition of the largest amounts of chromium for spontaneous passivation. In this connection, the passivating ability and, hence, the corrosion resistance based on spontaneous passivation sometimes increase when an environment becomes more aggressive. For instance, sputter-deposited W–Nb alloys are spontaneously passive in both 6M and 12M HCl, forming significantly tungsten-enriched double oxyhydroxide films of W4⫹ and Nb5⫹ ions, and the corrosion rates of the alloys with less than 52 at% niobium in 12M HCl are significantly lower than those in 6M HCl (33). Faster dissolution of niobium in 12M HCl leads to faster formation of the stable and protective passive films than in 6M HCl; hence, the corrosion damage in 12M HCl is significantly less than that in 6M HCl. 4. Effects of Metalloids As can be seen in Fig. 1, the corrosion resistance of amorphous alloys changes with additive metalloids, and the beneficial effect of metalloid in enhancing the corrosion resistance based on passivation decreases in the order phosphorus, carbon, silicon, and boron, as shown in Fig. 8 (34). The effect of metalloids on the corrosion resistance of alloys is dependent on the stability of the polyoxyanions contained in their films. Phosphorus and carbon contained in iron–chromium– metalloid alloys do not generally constitute passive films, as do phosphate and carbonate in strong acids, and they do not interfere with the formation of the passive hydrated chromium oxyhydroxide film (18,35). By contrast, as can be seen in Fig. 1, boron-containing alloys require larger amounts of the chromium Amorphous and Nanocrystalline Alloys 469 Fig. 8 Average corrosion rates of amorphous Fe–10 Cr–10 B–7 X and Fe–10 Cr–13 P–7 X alloys in 0.1N H 2SO 4 at 30°C. X is the metalloid denoted in the abscissa. (From Ref. 34.) addition in comparison with phosphorus-containing alloys in order to increase the passivating ability by the accumulation of chromium oxyhydroxide in the surface films because the films contain chromium borate (19,35). Silicon contained in amorphous metal–metalloid alloys forms surface films. Sputter-deposited Fe–Si alloys containing 25 at% or more silicon are passivated by anodic polarization due to the formation of a SiO2 film in a dilute sulfuric acid (36). The melt-spun amorphous Fe–39 Ni–10 B–12 Si alloy is more resistant to pitting corrosion in comparison with the amorphous Fe–40 Ni–20 B alloy because of the formation of a silicon-enriched surface film (37). An increase in silicon content of amorphous Fe–B–Si alloys extends the passive potential range 470 Hashimoto Fig. 9 Potentiodynamic polarization curves of amorphous Fe–10 Cr–5 Mo–B–Si alloys measured in 6M HCl at 25°C. (From Ref. 39.) (38). A special feature of silicon addition can be seen in Fig. 9 (39). Increasing the silicon content of amorphous Fe–10 Cr–5 Mo–B–Si alloys leads to a decrease in current densities in both the active and passive regions in 6M HCl at 25°C without changing the open-circuit corrosion potential. This is due to the formation of a SiO2-like substance covering the alloy surface along with the hydrated oxyhydroxide film, and the amount of the SiO2-like substance increases with increasing silicon content of the alloys. 5. Beneficial Effect of Phosphorus in Amorphous Alloys Phosphorus is well known to segregate at grain boundaries of stainless steels and to induce intergranular corrosion and stress-corrosion cracking of the stainless steels. However, phosphorus contained uniformly in amorphous alloys is effective in decreasing the corrosion rate and in enhancing passivation of amorphous alloys containing a passivating element. In addition to a large number of results including Figs. 1–5 and 8, another example can be given. When a variety of amorphous Fe–Cr–P alloys were prepared by ion beam mixing of a vacuum- Amorphous and Nanocrystalline Alloys 471 evaporaled iron, chromium, and iron phosphide multilayer, they showed remarkable corrosion resistance to sulfuric acids with and without chloride ions (40). The corrosion rate of amorphous Ni–P alloys is significantly lower than that of crystalline nickel, as shown in Fig. 5. Kinetic and analytical investigations (41) reveal the role of phosphorus as follows: Immersion of the amorphous Ni– P alloys in strong acids gives rise to preferential dissolution of nickel, leaving elemental phosphorus on the alloy surface. The elemental phosphorus is accumulated and forms the surface layer, which can act as the barrier layer to diffusion of nickel through the layer to dissolve into the solution. Accordingly, dissolution of nickel leads to thickening of the elemental phosphorus barrier layer and to the decrease of the dissolution current. In this situation, the decrease of the dissolution current with time at a constant potential follows Fick’s second law. A typical example is shown in Fig. 10. The equation in the figure is based on Fick’s second law, and the reciprocal of the current density squared changes linearly with polarization time. Phosphorus has another beneficial effect in enhancing corrosion resistance. Chromium contained in an alloy dissolves in the form of Cr 2⫹ ions in acids without oxidizing power, and if the oxidizing power is high, chromium is oxidized to Cr 3⫹ ions. Cr 3⫹ ions constitute a stable solid film with O 2⫺ and OH⫺ ions even in strong acids in the form of CrO x (OH)3⫺2x . This is the passive film of chromium. Fig. 10 Change in the reciprocal of the current density squared as a function of time of polarization at 100 mV versus saturated calomel electrode (SCE), where it is slightly anodic from the open-circuit potential. (From Ref. 41.) 472 Hashimoto Under natural conditions where no electricity is supplied by the outer circuit, the higher oxidizing power corresponds to the faster cathodic reactions, particularly oxygen reduction. The reduction rate of dissolved oxygen is mostly dependent on the activity of the surface for the oxygen reduction. The elemental phosphorus layer is especially active for the oxygen and proton reduction (42). The beneficial effect of metalloids in enhancing spontaneous passivation of amorphous and nanocrystalline Cr–P and Cr–B alloys due to acceleration of cathodic proton reduction is also pointed out by Moffat et al. (43,44). Accordingly, for amorphous alloys containing phosphorus and a strongly passivating element such as chromium, the formation of the elemental phosphorus layer not only prevents alloy dissolution acting as the diffusion barrier but also accelerates the passive film formation; that is, spontaneous passivation owing to the high activity for cathodic reactions. The beneficial effect of phosphorus in enhancing passivation in strong acids has been confirmed not only for Ni–Cr–P alloys but also for Fe–Cr–metalloid alloys (45) and Ni–Ta–P alloys (46). It has been found from these investigations that phosphorus-containing phases and/or metal phosphides are generally corrosion resistant even in strong acids and active for cathodic reductions. An interesting fact has been found in the role of phosphorus in the corrosion resistance of crystalline Cr–P alloys in a 47% HF solution at 30°C in which the concentration of dissolved oxygen is considerably higher than that in hydrochloric acids (47). The Cr–P alloys containing 0.7 at% or less phosphorus exhibit higher corrosion rates than that of chromium in the HF solution showing slightly higher open-circuit potentials than chromium in the active region of chromium, but the alloys with 0.8 at% or more phosphorus are spontaneously passive showing more than four orders of magnitude lower corrosion rates than chromium. The addition of small amounts of phosphorus enhances the cathodic oxygen reduction and hence accelerates active dissolution of the bcc chromium matrix phase. Accordingly, extended immersion results in complete dissolution of matrix leaving butterflylike flaky Cr 3P residues which previously exist in grain boundaries. This is the detrimental role of phosphorus in inducing intergranular corrosion and stress-corrosion cracking of crystalline alloys. By contrast, an increase in the phosphorus addition to Cr–P alloys leads to further enhancement of cathodic activity and to passivation of the matrix. C. Recent Efforts in Tailoring Extremely CorrosionResistant Alloys One serious restriction for the practical use of corrosion-resistant amorphous alloys prepared by melt-spinning methods is their limited thickness of several tens of microns, because the thickness of the melt which can be rapidly quenched from the liquid state for the formation of the amorphous structure is limited. Furthermore, conventional welding techniques cannot be applied, because heat- Amorphous and Nanocrystalline Alloys 473 ing of amorphous alloys leads to crystallization and to a loss of their superior characteristics based on the amorphous structure. One of the solutions to this problem is the preparation of amorphous surface alloys having the specific characteristics on conventional crystalline bulk metals. Tailoring new corrosion-resistant surface alloys has recently been performed mostly by sputter-deposition techniques. Sputter deposition is known to form a single amorphous phase alloy in the widest composition range among various methods for amorphization. 1. Aluminum-Corrosion-Resistant Metal Alloys Aluminum is the most widely used metal next to iron. Aluminum is not highly corrosion resistant and corroded in both acidic and alkali environments. Alloying with refractory metals, such as niobium, tantalum, molybdenum, and tungsten, is a potential method for enhancing corrosion resistance. However, the boiling point of aluminum is generally lower than the melting points of refractory metals. Accordingly, the conventional casting methods cannot be applied to the preparation of aluminum-refractory metal alloys. This bottleneck has been overcome by utilizing sputter-deposition techniques for alloy preparation. The sputter-deposition technique does not rely on melting to mix the alloying constituents and, hence, is suitable for forming a single-phase solid solution even when the boiling point of one component is lower than the melting point of the other components and/or when one component is immiscible with another component in the liquid state. The sputter-deposition method has been applied in tailoring corrosion-resistant aluminum alloys (48–54). Figure 11 shows the structure of various binary aluminum alloys (55). These alloys have been successfully prepared in a single amorphous phase in wide composition ranges. Alloying is very useful in enhancing the corrosion resistance. Figure 12 shows a comparison of corrosion rates of various aluminum alloys with those of corrosion-resistant conventional alloys measured in 1M HCl at 30°C. The use of aluminum alloys in 1M HCl has never been considered. However, when aluminum is alloyed with various corrosionresistant metals, the alloys possess sufficiently high corrosion resistance even in 1M HCl. Except for Al–Ti alloys, the corrosion resistance in 1M HCl increases with increasing alloying additions. Al–Ti and Al–Cr alloys dissolve actively, but other amorphous aluminum alloys are spontaneously passive even in 1M HCl. Amorphous Al–Ta and Al–Nb alloys are especially corrosion resistant. Sputter-deposition techniques have been widely used for the preparation of corrosion-resistant aluminum alloys in the first half of the 1990s. Shaw and her co-workers have found enhanced passivity of Al–W (56,57) and Al–Ta (58,59) alloys over a wide pH range due to the formation of tungsten and tantalum hydroxide films, and they have interpreted the high pitting resistance in terms of 474 Fig. 11 Hashimoto Structure of sputter-deposited aluminum alloys. (From Ref. 55.) an enrichment of alloying elements with high passivating ability in the underlying alloy surface. High pitting potentials of Al–V, Al–Mn, and Al–W alloys are attributed to the suppression of pit growth (60). Improved pitting resistance of Al–Ta alloys is attributed to the formation of thin passive films capable of impeding migration of chloride ions through the film (61), and when Al3Ta is precipitated, passivity breakdown occurs in the dealloyed region of the periphery of the cathodic precipitates (62). Amorphous Al–(15–45)Mo alloys show stable passivity over a wide potential range in 0.1M–1M NaCl, and their pitting potentials are 1.2 V or more positive than that of aluminum metal (63). In addition to sputter-deposited alloys, electrodeposited Al–Mn alloys form amorphous singlephase solid solutions and show considerably high pitting potentials (64). In this manner, the formation of single-phase solid solutions of aluminum alloys with corrosion-resistant elements is considerably effective in enhancing the corrosion resistance. 2. Chromium-Refractory Metal Alloys The valve metals such as titanium, zirconium, niobium, and tantalum are all passivated in strong acids. Chromium also has a very strong passivating ability. Consequently, if one succeeds in preparing chromium alloys with these valve metals, they seem to be ideal corrosion-resistant alloys in aqueous environments. These alloys have been successfully prepared by sputtering (65–70). Figure 13 shows the structure of sputter-deposited chromium–valve metal alloys as a function of alloying additions (66). These alloys form an amorphous single-phase Amorphous and Nanocrystalline Alloys 475 Fig. 12 Corrosion rates of various aluminum alloys and conventional corrosion-resistant alloys in 1M HCl at 30°C. (From Ref. 55.) Fig. 13 Structure of sputter-deposited chromium alloys. (From Ref. 66.) 476 Hashimoto solid-solution structure over wide composition ranges. They were all new amorphous alloys. Their corrosion resistance to concentrated hydrochloric acids is remarkably high. Figure 14 shows the change in the corrosion rates of Cr–Ti and Cr–Zr alloys in 6M HCl solution at 30°C and Cr–Nb and Cr–Ta alloys in 12M HCl solution at 30°C as a function of the valve-metal content of the alloys (65–68). In a 6M HCl solution, chromium and titanium dissolve actively, whereas the Cr– Ti alloys show very low corrosion rates, which are several orders of magnitude lower than those of the alloy components. Binary Cr–Zr alloys also show low corrosion rates. In spite of the fact that the corrosion rate of chromium metal is five orders of magnitude higher than that of zirconium metal, the corrosion rate of the Cr–Zr alloys decreases with increasing chromium content of the alloy. Amorphous Cr–Nb and Cr–Ta alloys show very high corrosion resistance, which is higher than that of the alloy components. These results indicate that if both components of binary alloys have a strong passivating ability, the alloys are able to possess better corrosion resistance than the alloy components. The corrosion rate of Cr–Ti, Cr–Zr, and Cr–Nb alloys tends to decrease with chromium content. Fig. 14 Corrosion rates of sputter-deposited Cr–Ti (65) and Cr–Zr (67) alloys in 6M HCl at 30°C and Cr–Nb (68) and Cr–Ta (68) alloys in 12M HCl at 30°C. Amorphous and Nanocrystalline Alloys 477 The corrosion rates of Cr–Ta alloys are extremely low and are lower than the level measurable by inductively coupled plasma (ICP) spectrometry. It can, therefore, be said that the amorphous Cr–Ta alloys are immune to corrosion even in 12M HCl and that the Cr–Ta alloys possess the highest corrosion resistance among known all metallic materials in strong acids. The chromium–valve metal alloys are all passivated spontaneously, forming the passive film consisting of both chromium and valve-metal cations. An interesting fact has been found with regard to the binding energy of core electrons of elements constituting the passive film. Figure 15 shows the correlation of the binding energies of the Cr 3⫹ 2p3/2 and Ta 5⫹ 4f 7/2 electrons with the cationic fraction of tantalum in the spontaneously passivated film formed on the Cr–Ta alloys (69). Their binding energies change with film composition, indicating the electronic interaction (i.e., charge transfer from chromium ion to tantalum ion in the film). Similar charge transfer from chromic ions, Cr 3⫹, to valve-metal cations has been found for Cr–Ti (65), Cr–Zr (67), and Cr–Nb (69) Fig. 15 Change in binding energies of the Cr 3⫹ 2p3/2 and Ta 5⫹ 4f 7/2 electrons with the cationic fraction of tantalum in the spontaneously passivated films formed on the Cr–Ta alloys in 12M HCl at 30°C. (From Ref. 69.) 478 Hashimoto alloys. The charge transfer between different cations indicates that these cations are located very closely to show the electronic interaction. This means that the passive film does not consist of a simple mixture of chromium oxyhydroxide and valve-metal oxide, but is composed of a double oxyhydroxide of chromium and valve-metal cations. The resultant double oxyhydroxide films are more protective than chromium oxyhydroxide and valve-metal oxide films in these aggressive solutions, and increasing chromium content of the alloys increases the chromium content of the double oxyhydroxide film and increases the corrosion resistance of the binary chromium–valve metal alloys. 3. Molybdenum-Corrosion-Resistant Metal Alloys Figure 16 shows corrosion rates of sputter-deposited molybdenum–valve metal alloys in 12M HCl (71–75). Molybdenum–zirconium alloys become amorphous in a wide composition range, whereas other molybdenum–valve metal alloys are Fig. 16 Corrosion rates of sputter-deposited molybdenum–corrosion-resistant metal alloys in 12M HCl at 30°C. (Data from Refs. 71–75.) Amorphous and Nanocrystalline Alloys 479 composed of nanocrystalline single bcc phases. Because grain diameters estimated from the full width at half-maxima of x-ray diffraction lines are 5–7 nm, these nanocrystalline alloys are regarded as homogeneous solid solutions from a corrosion point of view. All binary molybdenum–valve metal alloys show significantly higher corrosion resistance than that of alloy component elements regardless of crystalline and amorphous structures of the alloys. The corrosion rates of titanium, zirconium, and niobium are several orders of magnitude higher than the corrosion rate of molybdenum, but the corrosion rate of alloys decreases with increasing valve-metal content of the alloy. However, the corrosion rate of binary Cr–Mo alloys decreases with an increase in the molybdenum content of the alloy and the corrosion resistance never exceeds the corrosion resistance of molybdenum (75). The high corrosion resistance of the molybdenum alloys is attributable to the formation of passive double oxyhydroxide films of Mo 4⫹ and cations of alloying elements. Molybdenum alloys have different characteristics than chromium– valve metal alloys in film compositions and structures. Figure 17 shows polarization curves of Mo–Nb (73) and Mo–Ta (72) alloys measured in 12M HCl. All Fig. 17 Potentiodynamic polarization curves of Mo–Nb (73) and Mo–Ta (72) alloys measured in 12M HCl at 30°C. 480 Hashimoto polarization curves of molybdenum–valve metal alloys and Mo–Cr alloys measured in concentrated hydrochloric acids are similar to those shown in Fig. 17, although high zirconium alloys suffer pitting by anodic polarization at about 1.5 V (SCE) in 12M HCl. These molybdenum–corrosion-resistant element alloys are spontaneously passive and their open-circuit potentials are close to or higher than the open-circuit potential of molybdenum. Molybdenum dissolves actively from about ⫺0.8 to ⫺0.2 V (SCE) and passivates from about ⫺0.2 to 0.2 V (SCE), forming the film consisting of Mo4⫹ (76). However, molybdenum suffers transpassive dissolution at further higher potentials. As can be seen in Fig. 17, the cathodic activity of passive molybdenum for both proton and oxygen reductions is very high. Accordingly, the open-circuit potential of molybdenum in 12M HCl is very high and slight anodic polarization gives rise to a sharp current increase due to transpassive dissolution. Figure 18 shows cationic fractions of films and atomic fractions of underlying alloy surfaces of Mo–Nb (73) and Mo–Ta (72) alloys. Major anions in the films are O2⫺ ions and the balance is OH⫺ ions. The results shown in Fig. 18 are almost the same as those obtained for Mo–Ti (74), Mo–Zr (71), and Mo–Cr (75) Fig. 18 Cationic fractions of films and atomic fractions of underlying alloy surfaces of Mo–Nb (73) and Mo–Ta (72) alloys before and after immersion in 12M HCl at 30°C. Amorphous and Nanocrystalline Alloys 481 alloys either air-exposed or spontaneously passivated in concentrated hydrochloric acids. Therefore, both spontaneously passivated films and air-formed films are composed of oxyhydroxides containing cations of both alloy components. The air-formed films are generally rich in cations of alloying elements, because their affinity to oxygen is higher than that of molybdenum. Spontaneous passivation leads to the formation of molybdenum-enriched passive films on molybdenum-rich alloys and to the formation of the films with higher concentrations of alloying element cations on the alloys containing higher concentrations of alloying elements. There is a concentration gradient of cations in the films. Figure 19 shows the change in analytical results as a function of the photoelectron take-off angle (71,72). The measurement at a low photoelectron take-off angle emphasizes information from the exterior of the surface, and that at high photoelectron takeoff angle gives information not only from the exterior but also from the interior Fig. 19 Apparent cationic fractions of films and atomic fractions of underlying alloy surfaces of Mo–Zr (71) and Mo–Ta (72) alloys before and after immersion in 12M HCl at 30°C as a function of photoelectron take-off angle. 482 Hashimoto of the surface. Because of preferential oxidation of zirconium, the air-formed film on the Mo–Zr alloy is rich in zirconium, but immersion in 12M HCl requires an increase in molybdenum content in the interior of the film, which is necessary in avoiding corrosive dissolution. A similar increase in molybdenum content in the interior of the film is seen for the Mo–Ta alloy. Nevertheless, zirconium and tantalum cations are rich in the exterior of the passive films. This is generally found for Mo–Ti (74), Mo–Nb (73), and Mo–Cr (75) alloys. For spontaneous passivation, the air-formed film itself should be more or less stable and protective in 12M HCl. Although the film on the Mo–28 Ta alloy is rich in molybdenum on average, the increase in molybdenum content in the film by immersion in 12M HCl is not so high that the concentration of tantalum ions in the exterior of the film becomes lower than the tantalum content of the alloy. This is one of the reasons why the concentration gradient in the air-formed film on the Al–Ta alloy remains in the spontaneously passivated film. In addition, the open–circuit potential is located around the potential of transpassive dissolution of molybdenum; hence, molybdenum in the top-most surface of the film dissolves into the acid. Nevertheless, the high protective quality of the passive film is based on the synergistic effect of two cations forming the double oxyhydroxide film. Even if the alloy is polarized anodically over the transpassive potential of molybdenum, film-forming Mo 4⫹ ions are protected by chromium and valve-metal cations and are stable. Figure 20 shows changes in surface composition and oxidized molybdenum species for the Mo–57 Cr alloy polarized in 12M HCl as a function of polarization potential. Mo 6⫹ found in the film formed at potentials, where Mo 4⫹ is thermodynamically stable but Mo 6⫹ is not, results from air oxidation during washing in air after polarization and during transfer to the x-ray photoelectron spectrometer through air, as can be seen from the fact that molybdenum metal exposed to air is covered by Mo 6⫹ oxyhydroxide (e.g., Ref. 71). The content of Mo 6⫹ ions is almost constant in the potential region where Mo 4⫹ ions are major molybdenum ions and the Mo 6⫹ ions seem to be formed by air oxidation after polarization. The molybdenum species contributing to the protectiveness of the spontaneously passivated film is Mo 4⫹ ions, which are stable up to about 0.5 V (SCE) when the Mo 4⫹ ions are protected by chromic ions. When the polarization potential exceeds 0.6 V (SCE), protection by chromic ions is no longer effective and transpassivation of molybdenum occurs, showing a clear increase in the content of Mo6⫹ ions in the film. Consequently, the high corrosion resistance of sputter-deposited molybdenum alloys with other corrosion-resistant elements is due to the formation of double oxyhydroxide films consisting of Mo 4⫹ and cations of alloying elements, whose protectiveness is higher than oxyhydroxide films of alloy component metals. Amorphous and Nanocrystalline Alloys 483 Fig. 20 Changes in surface composition and oxidized molybdenum species for Mo–57 Cr alloy specimens polarized in 12M HCl as a function of polarization potential. 4. Tungsten-Corrosion-Resistant Metal Alloys Tungsten belongs to the same family as chromium and molybdenum in the periodic table; hence, it has been expected that tungsten–valve metal alloys are also extremely corrosion resistant. As shown in Fig. 21, four kinds of tungsten–valve metal alloys form amorphous single-phase structures (33,77–80). It is the first 484 Hashimoto Fig. 21 Structure of sputter-deposited tungsten–valve metal alloys. (Data from Refs. 33 and 77–80). time that W–Nb and W–Ta alloys were found to form single amorphous phase structures in spite of the complete miscibility of tungsten, with niobium or tantalum forming a continuous series of solid solutions at equilibrium. In general, if two alloy components form a continuous series of solid solutions at equilibrium, amorphous alloys are hardly formed, as can be seen in examples of Mo–Cr, Mo– Ti, Mo–Nb, Mo–Ta, and W–Cr. Furthermore, because sputter deposition is a very effective method for the formation of amorphous structure, when the composition is not suitable of forming the amorphous structure, their structure becomes nanocrystalline. Figure 22 shows corrosion rates of W–Ti alloys (77) in 6M HCl and W–Zr (78), W–Nb (33), W–Ta (80), and W–Cr (79) alloys in 12M HCl. In very aggressive hydrochloric acids, the binary alloys show lower corrosion rates than those of alloy component metals. The change in the corrosion rates of tungsten–valve metal alloys is very similar to that of molybdenum–valve metal alloys shown in Fig. 16. The corrosion rate decreases with increasing alloying elements and tantalum-containing alloys show particularly high corrosion resistance. W– Cr alloys also show higher corrosion resistance than tungsten and chromium, although the corrosion resistance of Mo–Cr alloys cannot exceed that of molybdenum. Figure 23 shows polarization curves of W–Zr alloys measured in 6M and 12M HCl at 30°C (78). Polarization curves of tungsten alloys are almost the same Amorphous and Nanocrystalline Alloys 485 Fig. 22 Corrosion rates of W–Ti alloys (77) in 6M HCl at 30°C and W–Zr (78), W– Nb (33), W–Ta (80), and W–Cr (79) alloys in 12M HCl at 30°C. as those of molybdenum alloys. However, a sharp continuous current increase due to transpassive dissolution as shown in Fig. 17 is not observed for tungsten alloys, but a sharp current increase with anodic polarization is followed by stagnation of current. The stagnant current of tungsten by anodic polarization is about 10 and 1 A/m2 in 12M and 6M HCl, respectively, and decreases with the addition of alloying elements (77–80). In this manner, the solubility of tungstate is not high and W6⫹ species remains in the film without transpassive dissolution. The spontaneously passivated films are composed of double oxyhydroxide of W4⫹ and cations of alloying elements. Because of no transpassive dissolution of tungsten, there is no clear concentration gradient in the depth of the films. This fact is similar to that found for chromium–valve metal alloys (81) and is a distinct difference from the passive films of molybdenum alloys in the exterior surface of which molybdenum is always deficient. 486 Hashimoto Fig. 23 Potentiodynamic polarization curves of W–Zr alloys measured in 6M and 12M HCl at 30°C. (From Ref. 78.) III. CORROSION-RESISTANT ALLOYS AT HIGH TEMPERATURES Another interesting fact is the extremely high resistance of aluminum-refractory metal alloys to high-temperature corrosion in sulfidizing and oxidizing environments. Corrosion of metallic materials in sulfur-containing atmospheres at high temperatures is much more severe than that in purely oxidizing environments. All conventional oxidation-resistant alloys suffer catastrophic corrosion in sulfurcontaining atmospheres at high temperatures, because of the poor protective properties of sulfide scales. For instance, the nonstoichiometry of sulfide scales formed on these alloys often reaches 10%. Because of the rapid diffusion of cations through the defective sulfide scale, they are sulfidized very rapidly. Some refractory metals such as molybdenum, niobium, and tantalum are, however, resistant to sulfide corrosion and their sulfidation rates are almost comparable to the oxidation rate of chromium. These metals, however, have very Amorphous and Nanocrystalline Alloys 487 low resistance to high-temperature oxidation in spite of the fact that practical sulfidizing atmospheres are often oxidizing. On the other hand, the best alloying element to form a protective scale in oxidizing environments is aluminum, and the second best is chromium. These metals form alumina and chromia scales, respectively. It has, therefore, been thought that aluminum-refractory metal alloys must be the best materials having high resistance to both oxidation and sulfidation. Sulfidation of sputter-deposited aluminum-refractory metal alloys, such as Al–Mo (82–85), Al–Nb (86–90), and Al–Ta (91) alloys follows a parabolic rate law, indicating that the rate-determining step of the overall reaction is the diffusional transport of matter through the sulfide scale formed. Figure 24 shows Fig. 24 Sulfidation (solid lines) and oxidation (dotted lines) rate constants for amorphous Al–Mo and Al–Mo–Si alloys as well as several high-temperature alloys. (Data from Refs. 82, 83, and 85.) 488 Hashimoto sulfidation (solid lines) and oxidation (dotted lines) rate constants for amorphous Al–Mo and Al–Mo–Si alloys as well as several high-temperature alloys (82,83,85). As can be seen from the comparison between solid and dotted lines, the sulfidation rate of conventional oxidation-resistant crystalline alloys is generally many orders of magnitude higher than the oxidation rate. By contrast, the sulfidation rates of Al–Mo and Al–Mo–Si alloys are significantly low and comparable to the oxidation rate of oxidation-resistant alloys. Furthermore, the sulfidation rate constants of these alloys are more than one order of magnitude lower than those of molybdenum, and the molybdenum-containing Fe–30 Mo–9 Al alloy. This result is of particular importance because this is the first time in the history of corrosion science that a metallic material has been obtained with a corrosion rate in highly sulfidizing atmosphere comparable to the oxidation rate of oxidation-resistant alloys. The steady-state sulfidation rates of Al–Nb (86– 90) and Al–Ta (91) alloys are also lower than those of the corresponding refractory metals. The sulfide scales on these alloys consist of two layers: the Al 2S3 outer layer and the inner refractory metal sulfide layer (84). The high sulfidation resistance of the Al–Mo alloys is attributed to the formation of the MoS 2 phase, constituting the major part of the inner barrier layer of the scale. The better protective properties of the sulfide scale formed on the Al–Mo alloys in comparison with those of the MoS 2 scale on molybdenum is attributed to a lower defect concentration in the aluminum-doped MoS 2 phase. The oxidation rate of Al–Mo alloys is comparable to that of chromia-forming alloys, although it is higher than that of alumina-forming alloys. However, the oxidation rate at temperatures higher than 900°C is very high. The scale consists mostly of alumina, but because of the high molybdenum contents of the alloys, molybdenum is also oxidized, forming volatile MoO 3. Because the melting point of MoO 3 is 793°C, the formation of low-melting point MoO 3 is responsible for the relatively low oxidation resistance of the Al–Mo alloys. Accordingly, an attempt to improve the oxidation resistance has been made by adding silicon to the Al–Mo alloys (83,85). The ternary Al–Mo–Si alloys have a high sulfidation resistance similar to that of the Al–Mo alloys and have a higher oxidation resistance than Al–Mo alloys. This is attributable to the formation of molybdenum silicide, which is stable to oxidation. During sulfidation and oxidation, amorphous alloys are crystallized, forming intermetallics. Al–Mo alloys form Al 8Mo 3 and Mo 3Al phases. The molybdenum-rich Mo 3Al phase is readily oxidized, forming volatile MoO 3. Accordingly, when the alumina scale surface on the Al–Mo alloys is analyzed, a low concentration of molybdenum is always found. By contrast, Al–Mo–Si alloys are crystallized to Al 8Mo 3 and Mo 5Si 3 phases without forming the easily oxidizable molybdenum-rich Mo 3Al phase. The Mo 5Si 3 phase is very stable to oxidation. Accordingly, any molybdenum and silicon are not detected Amorphous and Nanocrystalline Alloys 489 in the topmost surface of the alumina scale. The oxidation resistance of Al–Nb and Al–Ta alloys is also improved by the silicon addition (90,92). The chromium-refractory metal alloys also have a high sulfidation resistance, and the sulfidation rates of the alloys containing 50 at% or more refractory metals are almost comparable to those of corresponding refractory metals (93). The sulfide scales formed on these alloys consist of two layers: the Cr 2S 3 outer layer and the inner refractory metal sulfide layer. Cr–Nb and Cr–Ta alloys possess high oxidation resistance nearly comparable to typical chromia-forming alloys, although Cr–Mo alloys show poor oxidation resistance due to the formation of volatile MoO 3. IV. PREPARATION OF CORROSION-RESISTANT AMORPHOUS ALLOY COATINGS The practical use of extremely high corrosion resistance based on amorphous and nanocrystalline structures requires special methods for preparation. Laser and electron beam processing and sputter deposition are effective methods for the preparation of corrosion-resistant amorphous and nanocrystalline surface alloys. A. Laser and Electron Beam Processing Instantaneous irradiation of a metal surface by a high-energy-density beam such as a laser or electron beam is able to melt a small volume of the metal instantaneously. The heat of the melt can be rapidly absorbed by the large volume of cold solid metal surrounding the melt with a consequent rapid quenching of the melt. Figure 25 shows schematic drawings of laser and electron beam processing. If the surface composition is suitable, amorphization of a single irradiation trace of laser or electron beam has been found by many investigators. However, amorphization of a large surface area requires overlapping traverses by the high-energy-density beam and, hence, requires heating of the previously amorphized phase for irradiation melting of a portion adjacent to the previously amorphized phase. Thus, the crystalline phase appears in the heat-affected zone at the border of the neighboring irradiation traces of the high-energy-density beam (94). The prevention of crystallization in the previously amorphized zone during processing by the high-energy-density beam can be affected by proper selection of alloy composition having a higher glass-forming ability or by shortening of irradiation time by using a higher energy density. The first success in preparing corrosion-resistant amorphous surface alloys by laser treatment was made for Ni–Cr–P–B alloys on a mild steel substrate (95). At first, several tens of microns thick ribbon-shaped crystalline Ni–Cr–P– 490 Hashimoto Fig. 25 Schematic drawings of laser (a) and electron beam (b) processing of an amorphous surface alloy on a conventional metal substrate. B surface alloys were bonded to a bulk mild steel substrate by instantaneous furnace-melting the ribbons in vacuum and subsequent oil-quenching in vacuum. The previous vacuum melting was made for the formation of an adhesive bond between the crystalline surface alloy and the substrate metal, which was necessary for heat absorption from the melt to the substrate metal during laser treatment. Processing was performed by a continuous CO 2 laser beam and resulted in meltquenching at a thinner depth than the whole thickness of the surface alloy. Although Ni–Cr–P–B alloys containing 0–63 at% chromium became amorphous by the melt-spinning method, only Ni–(14–17)Cr–16 P–4 B alloys were completely amorphized by laser processing, and spontaneously passivated in 1M HCl. In this manner, the completely amorphizable composition range for the formation of the extremely corrosion-resistant amorphous surface alloys is seriously restricted. As far as melt-quenching methods are applied, the extremely high corrosion resistance is obtained only when the alloy becomes completely amorphous. Because melt-spun amorphous nickel–valve-metal–platinum group metal alloys have very high activity and durability for the anode of seawater electrolysis to form sodium hypochlorite, corrosion-resistant wide electrode plates covered with amorphous surface alloys were prepared by high-energy-density beam processing (96). At first, nickel was electrodeposited on a niobium substrate and then a salt of a platinum group element was coated on the surface of the nickel. Then, the three-layered specimen was heat treated to form an adhesive bond between the nickel and niobium. Although Ni–(25–65) Nb alloys were amorphizable by melt-spinning, only Ni–(35–40)Nb alloys were amorphized over the Amorphous and Nanocrystalline Alloys 491 whole surfaces by laser processing. Unless complete amorphization of the whole surface occurred, in addition to crystallization in the heat-affected zone of neighboring laser traces, a number of cracks were formed as a result of solidification of melt under restriction by the solid substrate. The laser processing is rather slow, mostly because of a high reflectivity of the infrared CO 2 laser beam of 10.6 µm wavelength from solid surfaces of metals such as nickel. On the other hand, an electron beam is another high-energy-density beam and is easily absorbed by metals. Accordingly, the amorphous Ni–Nb–platinum group metal surface alloys on niobium were prepared by both laser and electron beam processing (97). It was found that only 1.7% of the CO 2 laser beam is absorbed by the metal specimen when the electron beam is assumed to be absorbed by the metal specimen with 100% efficiency (98). For instance, when a 6-kW CO 2 laser and electron beam machines are used for the formation of a 1-m 2 amorphous surface alloy from a 15-µm-thick nickel-plated niobium specimen, the processing times are 7.4 h and 22 min, respectively. Consequently, for the preparation of a plane amorphous surface alloy, electron beam processing is more convenient than laser processing. It was found (99,100) that the energy consumption for electrolysis of seawater by the corrosion-resistant amorphous surface alloy anodes was about two-thirds of that by the currently used most active electrode. In this manner, laser and electron beam processing is one of effective methods for the preparation of amorphous surface alloys. However, because amorphizable compositions are especially limited and because manufacturing processes are rather complicated, laser and electron beam processing is only suitable in producing high-value-added special materials which cannot be prepared by any other methods. B. Sputter-Deposition Coating of Inner Wall of Tube As mentioned earlier, sputter deposition is one of the best methods to form amorphous and nanocrystalline surface alloys. In general, conventional sputtering methods can be applied for the formation of plane surface alloys. The practical application of a corrosion-resistant amorphous alloy to various plants requires sputter-coating on inner walls of narrow tubes. This has been performed by Shimamura (101). Figure 26 shows a schematic drawing of the sputtering machine designed by Shimamura. A water-cooled thin target rod is inserted into a substrate tube and magnetron sputtering can be done using an electric magnet placed outside the sputtering chamber. On the left-hand side of the sputtering chamber, previous cleaning of target can be done by radio-frequency sputtering. He used a target consisting of Type 316L stainless steel, on the surface of which four tantalum bars are embedded symmetrically. Accordingly, sputter deposition of amorphous Fe–Ni–Cr–Ta alloys on the inner walls of Type 304LTP stainless 492 Hashimoto Fig. 26 A schematic drawing of a sputter-deposition machine for coating of the inner wall of a narrow tube and the cross section of the target (From Ref. 101.) steel tubes of 38.4 and 54.9 mm inner diameter was performed. The sputter deposits showed good performance in boiling 8M HNO 3 solutions containing various oxidizing cations such as Cr 6⫹. V. CONCLUDING REMARKS Almost any kinds of properties can be obtained by the preparation of amorphous and nanocrystalline alloys, mostly because of the formation of homogeneous alloys due to the expansion of solubility limits. A variety of corrosion-resistant materials is able to be tailored depending on the aim and environmental conditions. Bulk amorphous alloys can be processed if the alloys have a wide gap between glass transition and crystallization temperatures, but their compositions are restricted. 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Laser and electron beam processing of amorphous surface alloys on conventional crystalline metals. Mater Manuf Processes 5:567–590, 1990. K Shimamura, DEng dissertation, Tohoku University, Sendai, 1996. Index Acids (see Environments for corrosion) Aerospace, 375 Agents (fiber-reinforced materials): mold release, 421 silane coupling, 421 Aircraft (fiber-reinforced materials), 420 Alkalis (see Environments for corrosion) Alloys, amorphous and nanocrystalline: behavior at high temperatures, 486– 489 high corrosion resistance, 460 coatings of amorphous alloys, 489–492 factors determining resistance, 461–472 tailoring corrosion-resistant alloys, 472–486 Aluminum, 173–251 alloys: Al-Cu, 216–220 Al-Li, 225 Al-Mg, 220–222 Al-Zn, 220–223 amorphous, 473, 474 commercial, 230–232 ternary, 230–232 applications, 173, 174 [Aluminum] corrosion (see Corrosion and Environments for corrosion) Amorphous (see Alloys) Anneal (heat treatment): amorphous alloys, 462–465 ceramic matrix composites, 412 copper, 130, 139 iron, 23 FeAl, 289 magnesium, 254–258, 264, 270 metal matrix composites, 380 nickel, 61, 103, 104 stainless steel, 45 Antenna, 375 Atmosphere (see Environments for Corrosion) Austenitic (see Steel, types of) Automotive: fiber-reinforced materials, 420 metal matrix composites, 375 Bainitic (see Steel, types of) Bases (see Environments for corrosion) Battery, high temperature, 103 Bearings, ball, 314 Blisters (formation on aluminum), 205– 207 501 502 Boiler: coal or oil fired, 88, 100 waste heat, 91 Boron (suppresses environmental embrittlement), 277–283 Brass: single-phase alloy, 118 two-phase alloy, 119 Bridge (fiber-reinforced materials), 419, 420 Bridging, crack wake, 391 Buildings (fiber-reinforced materials), 419, 420 Cavitation, 126 Cell, fuel, 103 Cement (see Environments for corrosion) Ceramics, BN, AIN, silica forming, 316–323 Ceramics, non-oxide, 311–344: corrosion (see Environments for corrosion) effects of oxidation/corrosion, 342– 344 oxygen transport in silica, 315, 316 Ceramics, oxide, 351–416: corrosion (see Corrosion and Environments for corrosion) Chlorination, 95–99 Chromium (amorphous alloys), 474– 478 Coatings (see Corrosion, mitigation) Composites: ceramic matrix, 391, 392 glass, 413–415 oxides, 410–413 silicon carbide, 392–410 metal matrix, 375 aluminum, 376–381, 384, 385 iron, 386 magnesium, 378, 379, 381, 382, 385, 386 titanium, 382, 386 metal matrix corrosion (see Corrosion) Index Composition: grain-boundary (steel), 41–44 equilibrium processes, 42, 44 nonequilibrium processes, 45 Fe3Al, 294 solution (see Environments for corrosion) Concrete (see Environments for corrosion) Conductivity, solution, 48, 49 Containment, 103, 109 Contamination (see Environments for corrosion) Copper: corrosion (see Corrosion and Environments for corrosion) general, 115, 116 Corrosion: causes of (see Environments for corrosion) cracking (see stress corrosion cracking) crevice corrosion: aluminum, 177, 234, 238 copper, 118 metal matrix, 379 steel, 6 titanium (saltwater), 156 dealloying: copper, 118, 119 steel, 6 erosion-corrosion: aluminum, 243–245 copper, 125 steel, 7 exfoliation of aluminum, 178, 180 fatigue: aluminum, 180 copper, 121–123 FeAl (crack growth), 292 Fe3Al (crack growth), 294, 295 magnesium, 267–270, 385 plastics, fiber reinforced, 433 reinforced aluminum, 384 static (oxide ceramics), 356–359 steel, 6 titanium (crack growth), 386 Index [Corrosion:] galvanic: aluminum, 177, 178 copper, 126 magnesium, 272 metal matrix composites, 379, 380, 382 plastics, fiber reinforced, 422 steel, 7, 20, 21 titanium (saltwater), 156 general: aluminum, 174, 175 copper, 116, 117 magnesium, 272 steel, 2, 3, 5, 21–29, 45, 75–109 hydrogen-stress cracking (steel), 7, 10, 11, 16, 17 intergranular: aluminum, 175, 230, 241, 244 copper, 123–125 steel, 6, 32 localized (aluminum), 175 mechanisms, fiber-reinforced plastics: fiber based, 447 interfaced based, 448 matrix based, 448 multistage, 448, 449 pitting: aluminum, 175, 208–212, 221– 225, 228–231, 233 amorphous alloys, 473, 474 composites, metal matrix, 376 copper, 117, 118 magnesium, 265 metal matrix composites, 377, 379 steel, 6–10, 32 zirconium, 163, 166 problems, fiber-reinforced plastics: scale, 449 variability of materials, 449, 450 resistance, amorphous Fe-Cr-metalloid alloys, 460 stress corrosion cracking: aluminum, 180, 238, 241 composite materials, model of, 386–388 503 [Corrosion:] copper, 119–121 intermetallics, 285, 296, 383– 385 iron matrix composite, 386 magnesium, 266–271 plastics, fiber reinforced, 443–445 steel, 7, 10–19, 32, 41–44, 47, 48 Corrosion, mitigation: anodic and cathodic protection: amorphous alloys, 472, 480 FeAl, 290–291 metal matrix composites, 381, 388, 389 steel, 3, 5, 21 coatings: amorphous alloys, 489–492 on ceramics (refractory oxide), 339 on intermetallics, 304 magnesium, 271 on steel 3, 27–29 Cracking: hydrogen-stress (see Corrosion) matrix, 391 stress-corrosion (see Corrosion) Crevice (see Corrosion) Dealloying (see Corrosion) Debonding, interface, 391 Deflection, crack, 391 Dezincification, 118, 119, 132, 139 Discontinuities, microstructural (steel): grain boundaries, 9 second-phase particles, 9 sulfide inclusions, 9 Duplex (see Steel, types of ) Electron beam (producing amorphous alloys), 489–491 Electronics, high temperature, 311 Embrittlement: alleviation, 303–307 alloy additions (B, Cr, Zn), 306 grain size and shape, 304, 305 prestrain, 306 504 [Embrittlement:] processing techniques, 306 surface conditions, 303, 304 elevated temperatures, 298–303 L12 alloys (Ni3Al, Ni3Si), 299–301 other intermetallics, 301, 302 environmental, suppressed by: boron, 277–283 zirconium, 283 hydrogen induced, 296, 297 ceramic matrix composites, 408–410 intermetallics (see Intermetallics) metallic alloys, 162, 170, 172 titanium, 158 mechanisms, 302, 303 moisture induced, 284, 297, 307 oxygen induced, 307 Energy, solar, 103, 107 Engine: gas turbine, 414 heat, 311, 323 Environments, aqueous (aluminum), 193–232 alloys in water, 215, 216 binary alloys in water: Al-Cu, 216–220 Al-Mg, 216–221 Al-Zn, 220 unalloyed alloys in water, 195 adsorption process, 195–201 chemical reaction, 201–204 direct attack of exposed metal, 205–207 thinning of oxide film, 204, 205 Environments for corrosion: acids, effect on: amorphous metalloid alloys, 460 copper, 146, 147 plastics, fiber reinforced, 441, 447 steel, 40 titanium (oxidizing and reducing), 157 alkalis, effect on: copper, 147 plastics, fiber reinforced, 433, 441 titanium, 158 Index [Environments for corrosion:] ammoniacal solutions, effect on copper, 145, 146 antifreeze, effect on fiber-reinforced plastics, 440 aqueous, effect on aluminum (see Environments, aqueous) ash- and salt-deposits, effect on nickel (high temperature), 75, 76, 99–103 atmosphere, effect on: copper, 130–134 steel, 26, 38 base, effect on steel, 40 carbon dioxide, effect on silicon carbide, 329 carburization and metal dusting, effect on nickel (high temperature), 75, 76, 88–91 cement, effect on fiber-reinforced plastics, 423, 431, 441, 447 chloride concentration, effect on aluminum, 213 concrete, steel rebar, 26, 27 contamination, 25, 45 deionized water, effect on steel, 39 diesel fuel, effect on fiber-reinforced plastics, 440 environmental degradation, effect on fiber-reinforced plastics, 419, 453 erosion, flow assisted, effect on steel, 7 freshwater, effect on: copper, 139–141 steel, 25, 26 gas, effect on: fiber-reinforced plastics, 440 titanium, 158, 159 halogens, effect on: nickel (high temperature), 75, 76, 95–99 zirconium, 167, 168 inorganic acids, effect on titanium, 157, 158 lubricating oil, effect on fiber-reinforced plastics, 440 Index [Environments for corrosion:] melts, thick, effect on oxide ceramics, 371 metals, liquid (molten): effect on oxide ceramics, 370– 372 nickel (high temperature), 75, 76, 109 titanium, 159 neutral solutions, effect on copper, 147 nitrogen, effect on oxide ceramics, 355 nitridation, effect on nickel (high temperature), 75, 76, 91–95 organic compounds and solutions, effect on: copper, 147 titanium, 158 oxidation, effect on: aluminum, 221, 232, 233 (See Films, surface) amorphous alloys (high temperature), 486–489 ceramic matrix composites, 413 ceramics, non-oxide, 314–325, 332–339 copper, 128–130, 146 magnesium, 265–266 nickel (high temperature), 75, 76– 85 silicon carbide, 398–408 steel, 2 oxygen, effect on oxide ceramics, 352–355 pH, effect on: aluminum, 214, 237 steel, 2, 25 pollution, effect on copper, 144, 145 potable water, effect on: copper, 137–139 steel, 39 pressure, hydrostatic, effect on fiberreinforced plastics, 445 radiation, effect on fiber-reinforced plastics, 442 505 [Environments for corrosion:] saline solution (saltwater, seawater), effect on: aluminum, 218, 222, 225, 231, 235, 241–245 copper, 141–144 magnesium, 261–265 plastics, fiber reinforced, 433 steel, 39, 40 titanium, 156, 157 salts, molten, effect on: nickel (high temperature), 75, 76, 103–109 oxide ceramics, 359–371 salt spray, effect on magnesium matrix, 378, 379 soil, effect on: copper, 136, 137 steel, 26, 27 solution composition, 2 steam, effect on copper, 145 sulfidation, effect on: amorphous alloys (high temperature), 486–489 nickel (high temperature), 75, 76, 85–88 temperature, effect on: aluminum, 214, 215 amorphous alloys, 486–489 ceramic matrix composites, 402– 407 ceramics, non-oxide, 339 glass matrix composites, 413–415 nickel, 75–109 plastics, fiber reinforced, 425, 426, 432, 440–442, 451 steel, 2 velocity, fluid, effect on steel, 2, 23, 24 water, effect on: aluminum (see Environments, aqueous) ceramics, non-oxide, 325–329, 340, 341 ceramics, oxide, 356–359 copper, 134–136 506 Index [Environments for corrosion:] magnesium, 258–2 70 plastics, fiber reinforced, 422, 423, 426, 432, 434, 440–442 titanium, 155, 156 Epoxy, 422 Erosion-Corrosion (see Environments for corrosion) Exchanger, heat, 414 Extraction, metallurgical, 103 Fumes, gasoline and motor oil, 423 Furnace: carburizing, 91 sulfur, 87 Fatigue (see Corrosion) Feedstocks, chemical, 102 Ferritic (see Steel, types of) Fibers: alumina, 385 aramid, 422, 436 boron, 379, 423 carbon (graphite), 380, 422, 436 continuous, 391 glass, 422, 435 metal, 423 plastics, structure of, 421 Fillers (fiber-reinforced materials), 421 Film, passive (steel), 31 Film, surface (aluminum), 181, 232– 245: dry environments, 181, 182 wet environments, 182–193, 207–209 aqueous environments (see Environments, aqueous) electrochemical potential, 207–209 40°C to 100°C, 183–185 100°C to 150°C, 185–190 ⬎150°C, 190–193 room temperature to 40°C, 183 Filter, porous, 311 Fins, vertical tail, 375 Formers: alumina, 84, 85, 89 chromia, 84, 85, 89 silica, effect of suboxide, 329–332 Fracture, low ductility and brittle, 275 Freshwater (see Environments for corrosion) Fuel, sulfur bearing, 88 Hafnium, 169 Hardening, precipitation (see Steel, types of) Hardware, military, 375 Heat treatment (see Anneal) Highways (fiber-reinforced materials), 419 Homogeneity (amorphous alloys), 462– 465 Humidity, relative (steel), 25 Hydrostatic (see Pressure, hydrostatic) Galvanic (see Corrosion) Gillotine (test method for aluminum), 205 Glasses, silicate, 425, 433, 434 Guidance, inertial, 375 Impurities, effect on non-oxide ceramics, 332–339 Incinerator, municipal waste, 101 Inhibitors to corrosion (steel), 3 Intergranular (see Corrosion) Intermetallics, environmental embrittlement, 275–307: Ll2, 275–286 Ni3Al, 277–285 Ni3Fe, 277 Ni3Si, 285, 286 iron aluminide: FeAl, 287–292 Fe3Al, 287–292 NiAl, 286–287 other (iron or nickel based), 296, 297 Interphase, fiber-reinforced plastics, 440 Ion-implantation (aluminum), 225, 226 Iron, cast, 1, 23, 62 Ketone, polyether (see Resins) Laser (producing amorphous alloys), 489 Index Liners: combustor, 311 fiber-reinforced material, 420 Lines, groundwater (copper), 115 Magnesium, 253–272: Mg-Al, 254–258 Mg-Zn, 254–258 Marine (see Vessels, marine) Martensitic (see Steel, types of ) Mechanisms, toughening, 391 Melt-spinning (aluminum), 225 Metalloids, effects on amorphous alloys, 468–470 Metals: liquid, 341, 342 reactive, 151, 152 refractory, 151, 152 Microelectronics, silicon, 318 Modes of corrosion (see Environments for corrosion) Molybdenum (amorphous alloys), 478– 483 Motors, rocket (fiber-reinforced materials), 420 Nanocrystalline (see Alloys) Nickel: advantages, 55, 56 corrosion (see Corrosion and Environments for corrosion) processing chemicals: chlorides, 71, 72 hydrochloric acid, 65–68 hydrofluoric acid, 68, 69 hydroxides, 72, 73 nitric acid, 70 phosphoric acid, 69 sulfuric acid, 63–65 resistance to high-temperature corrosion (see Environments for corrosion) systems: Ni, 56–58 Ni-Cr, 58 Ni-Cr-Si, 62 507 [Nickel:] Ni-Cu, 58–60 Ni-Fe-Cr, 63 Ni-Mo, 60, 61 Niobium, 170–171 Nonequilibrium (aluminum), 225 Optics, precision, 375 Organic compounds (see Environments for corrosion) Oxidation (see Environments for corrosion) Oxygen, dissolved (steel), 25 Particulate, 391 Passivation (amorphous alloys): high activity, 465–468 high passivating ability, 461–462 Pearlitic (see Steel, types of ) pH (see Environments for corrosion) Phenolic (see Resins) Phosphorus in amoprhous alloys, 470– 472 Pipelines (fiber-reinforced materials), 420 Pitting (see Corrosion) Plastics, fiber reinforced: corrosion (see Corrosion and Environments for corrosion) environments, 423 interphase degradation, 439–440 mechanical properties, 429–431 porosity, 450 predicting long-term performance, 451–453 structure, 421–423 Platforms (fiber-reinforced materials), 420 Poisoning (see Toxicity of cuprous oxide) Polarization (aluminum), 209–212 Polybismaleimides (see Resins) Polyester, 422 Polyether (see Resins, polyether ketone) Polyetherimide (see Resins) Polyethersulfone (see Resins) 508 Polyimides (see Resins) Polymers: degradation, 436 thermoplastic, 437 thermosetting, 421, 437 Polysulfone (see Resins) Potential: critical pitting (aluminum), 208, 209 protection (aluminum), 209 Pressure, hydrostatic (aluminum), 242, 243 Protection, passivation (steel): factors that hinder: concentration changes, 5 impurities, 5 temperature increase, 5 velocity changes, 5 factors that promote: hydrofluoric acid, 3 sodium hydroxide, 3 sulfuric acid, 3 Pullout, fiber, 391 Rail, light (fiber-reinforced materials), 420 Rainfall (see Environments for corrosion) Reactor, fusion, 171 Recuperators, 102 Reinforcement: ceramic matrix composites (fiber), 396, 397 metal matrix composites continuous fiber, 379–382, 384– 388 discontinuous, 376–379, 383–384 Repassivation (aluminum), 223, 230 Resins: phenolic, 422 polybismaleimides, 422 polyetherimide (thermoplastics), 421 polyether ketone (thermoplastics), 421 polyethersulfone (thermoplastics), 421 polyimides, 422 polysulfone (thermoplastics), 421 Index Resistance to corrosion (see Corrosion, mitigation) Roads (fiber-reinforced materials), 420 Rocket (fiber-reinforced materials), 420 Semiconductor, high temperature, 313 Shield, reentry (space vehicle), 311, 313 Silane (see Agents, silane coupling) Silicon nitride, oxidation of (see Ceramics, non-oxide) Silicon, oxidation of (see Ceramics, non-oxide) Silicon carbide, oxidation of (see Ceramics, non-oxide) Sizing (fiber-reinforced materials), 421 Soil (see Environments for corrosion) Solar (see Energy) Solidification, rapid (aluminum), 225 Solutions, neutral (copper) (see Environments for corrosion) Sputter-deposition: aluminum, 225, 226, 229, 230 amorphous alloys, 473, 491, 492 magnesium, 271 metal matrix composites, 389 Stairs (fiber-reinforced materials), 420 Steel: carbon, 2, 6 corrision (see Corrosion and Environments for corrosion) elements of stainless steel: carbon, 37 chromium, 36 manganese, 36 molybdenum, 36 nickel, 37 nitrogen, 37 low-alloy, 2 stainless steel, types of: austenitic, 31–33, 48–50 bainitic, 23 duplex, 32, 33, 51 ferritic, 1, 23, 32, 33, 50, 51 martensitic, 1, 23, 32, 33, 51, 52 pearlitic, 23 precipitation hardening, 32, 33, 36 Index [Steel:] strength: high-strength (see Corrosion, stress corrosion) low-strength (see Corrosion, stress corrosion) weathering, 22 Stiffeners, cargo bay, 375 Stress, threshold (steel), 13 Substrate, electronic, 313 Superaustenitics, 50 Superduplex, 51 Supermartensitics, 52 Tanks: chemical storage (fiber-reinforced materials), 420 oil storage (fiber-reinforced materials), 420 Tantalum, 171–172 Tarnishing, 135, 136 Temperature (see Environments for corrosion) Tests: accelerated, 424, 425 service condition, 424 Titanium, 154–159: applications, 154, 155 alloys (alpha, alpha-beta, beta), 155 509 [Titanium:] corrosion resistance (see Environments for corrosion) Toxicity of cuprous oxide, 144 Tubes: copper: condenser, 115 plumbing, 115 heat exchanger, 311 Tungsten (amorphous alloys), 483–485 Vanadium, 169–170 Velocity, fluid (see Environments for corrosion) Vessels: marine (fiber-reinforced materials), 420 reaction (fiber-reinforced materials), 420 Vinylester, 421, 422 Water (see Environments for corrosion) Whiskers, 376, 391, 396 Zirconium: general, 159–170 nuclear applications, 160, 161, 167 suppresses environmental embrittlement, 283
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