Development of new Co–Cr–W-based biomedical alloys: Effectsof microalloying and thermomechanical processing on microstructures and mechanical properties Kenta Yamanaka a,⇑ , Manami Mori b , Koji Kuramoto a,c , Akihiko Chiba a a Institute for Materials Research (IMR), Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan b NISSAN ARC, LTD., 1 Natsushima-cho, Yokosuka 237-0061, Japan c Eiwa Co., Ltd., 405-45 Kasshi-cho, Kamaishi 026-0001, Japan a r t i c l e i n f o Article history: Received 10 August 2013 Accepted 18 October 2013 Available online 30 October 2013 Keywords: Biomedical Co–Cr–W alloy Thermomechanical processing Microstructure Precipitates Mechanical properties a b s t r a c t The application of computer-aided design and computer-aided manufacturing (CAD/CAM) to dentistry has recently attracted considerable attention as a new technique for designing and fabricating custom- made dental implants. Here, a strategy combining microalloying with thermomechanical processing are described to design new Co–28Cr–9W–1Si–C (wt%) alloys for use as disks in the CAD/CAM-based machining of dental restorations. On the basis of our thermodynamic calculations, Si and C were selected as alloying elements that cause the brittle r phase precipitates to be replaced with the plastically deformable Laves phase and thus enhance the alloy’s hot workability. The effect of thermomechanical processing on the microstructure evolution and mechanical properties of the designed alloys was prelim- inarily studied by performing multipass hot rolling. The hot-rolled alloys exhibited refined grains (mean grain sizes 10 lm) and high densities of lattice defects (dislocations, stacking faults, etc.), both of which were obtained as a result of dynamic recrystallization during hot rolling. It was found experimentally that this approach permits the alloy strength and ductility to be increased simultaneously. The static recrys- tallization occurring during cooling after deformation also modifies the mechanical properties of the alloys. Carbon doping (<0.1 wt%) increases the amount of precipitates and further improves both the strength and elongation-to-failure of the hot-rolled alloys. Thus, the newly developed alloys have advan- tageous characteristics in terms of both fabrication and mechanical properties. In addition, the outstand- ing tensile ductility of the developed alloys could make them suitable for vascular stents. Ó 2013 Elsevier Ltd. All rights reserved. 1. Introduction Co–Cr alloys with good corrosion and wear resistances have been used in various orthopedic implants such as artificial hip and knee joints, and they are also used in dentistry. Recently, com- puter-aided design and computer-aided manufacturing (CAD/ CAM) technologies have been applied as novel methods for design- ing and fabricating dental restorations such as crowns, bridges, and inlays. This strategy provides a rapid, low-cost, and precise means of fabricating custom-made dental restorations for patients. In particular, selective laser melting [1,2] and CAD/CAM-based machining [3–5] have been proposed in this context. In the CAD/ CAM-based machining process, products are milled from ceramic, composite resin, or metallic disks based on three-dimensional CAD data. The manufacturing of all-ceramic dental restorations consisting of zirconia-based materials is the most frequent applica- tion of this method [4]. On the other hand, metal-ceramic systems also play an important role because they can realize a good combi- nation of mechanical rigidity and aesthetics, originating from the metallic framework and ceramic veneer, respectively [5,6]. Although biocompatible gold-based alloys such as Au–Ag–Pd al- loys have usually been employed as the metallic components [7], Co–Cr or Ni–Cr alloys, which consist of nonprecious elements, are becoming increasingly popular for this application because of high cost of gold. It is also widely accepted that Ni may cause allergies and cancer in living organisms, and therefore Ni-contain- ing alloys are not suitable for use in biomaterials [8]. In addition, the Co–Cr-based alloys have higher corrosion resistance than Ni–Cr alloys [9]. Therefore, Ni-free Co–28Cr–9W-based (wt%) cast alloys with high biocompatibility have been commercialized. How- ever, these cast alloys generally have large grains and exhibit solid- ification segregation, which results in inadequate mechanical properties and inhomogeneous material quality. Although homog- enizing heat treatment can be used to eliminate this segregation, unavoidable grain growth still adversely affects the mechanical and milling properties of these alloys. 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.10.052 ⇑ Corresponding author. Tel.: +81 22 215 2118; fax: +81 22 215 2116. E-mail address:
[email protected] (K. Yamanaka). Materials and Design 55 (2014) 987–998 Contents lists available at ScienceDirect Materials and Design j our nal homepage: www. el sevi er . com/ l ocat e/ mat des Furthermore, Co–Cr-based dental alloys are required to have yield stresses higher than 500 MPa [10]. The higher strength basi- cally leads to higher fatigue strength and can improve the mechan- ical reliability of restorations that are subjected to occlusal forces. Therefore, in recent years, there has been extensive research on improving the mechanical properties of Co–Cr-based dental alloys [1,11,12]. In addition, materials used in these applications should consist of small grains, because chipping failure occurs in ma- chined components with coarse grain structures, reducing the accuracy of the fit of a restoration. Thus, a grain refinement process is necessary to improve these properties. The present study describes a strategy for designing a new class of Ni-free Co–Cr–W-based alloys that have uniform, refined micro- structures and exhibit excellent mechanical properties. To accom- plish this goal, hot deformation was performed to optimize their microstructures. Previously, it was found that bulk nanostructured Co–Cr–Mo alloys with average grain sizes of less than 1 lm can be obtained by hot deformation with very low applied strains (equiv- alent strain, e eq , <1) [13–17], although e eq over 4 is generally re- quired to realize such microstructures [18]. This significant grain refinement is derived from dynamic recrystallization (DRX), which is closely related to an extremely low stacking fault energy (SFE) at high deformation temperatures (>1273 K) [13,16]. In addition to the grain refinement, many more dislocations are accumulated during hot deformation of the Co–29Cr–6Mo alloys with very low SFEs than in conventional alloys [17]. Accordingly, the applica- tion of hot deformation dramatically improves their tensile proper- ties [14,15,17]. Although a similar approach would be a promising means of obtaining refined microstructures and enhanced mechanical properties in Ni-free Co–Cr–W alloys, no studies have been reported on hot deformation processing and its effects on their structure–mechanical property relationships. Therefore, as a preliminary study, multipass hot rolling of the Co–28Cr–9W-based alloys was conducted and its influence on their microstructures and mechanical behaviors were investigated. The present study also considered the alloying strategy, because the constituent phases influence the alloy’s mechanical characteristics. Before the experiments, the alloying elements were modified based on ther- modynamic calculations to improve their hot workability and room-temperature mechanical properties. 2. Alloy design The alloying strategy for developing novel Ni-free Co–Cr–W- based alloys is discussed below. Firstly, the effects of Cr and W on the microstructural evolution are summarized. Biomedical Co–Cr alloys generally have face-centered cubic (fcc) c and hexagonal close-packed (hcp) e phases. The addition of chromium stabilizes the e phase [19], while increasing the tung- sten content increases the c volume fraction [20]. A sufficiently high c-phase stability is desirable in order to obtain better hot workability relative to that of the Co–Cr–Mo alloys, because the plastic deformation occurring when the e phase exists leads to fractures [21]. Karaali et al. also reported that the addition of tung- sten resulted in precipitation of the r phase, which is an undesir- able brittle intermetallic compound (IMC) [20]. Our preliminary study revealed that hot-forged Co–Cr–W-based alloys with large amounts of tungsten (15 wt%) fail in the elastic regime during tensile testing (Kurosu S, Sugihara K, Matsumoto H, Koizumi Y, Chi- ba A. unpublished result). Then, Si and C were selected as the min- or alloying elements and examined their effects on the c ?e transformation and precipitation behavior. First, the effect of Si on the phase equilibria was investigated. Si (up to 2 wt%) has previously been used only to enhance the cast- ability of Co–Cr-based dental alloys [22]. Fig. 1a shows a vertical section of the phase diagram of the Co–28Cr–9W–Si system calcu- lated using Thermo-Calc (the thermodynamic data sets used in this calculation were obtained from TCS Steels/Fe-alloys Database Ver. 6). According to the calculated phase diagram, the fcc c phase is stable only at high temperatures, and the equilibrium matrix phase at room temperature is the hcp e phase. Si stabilizes the e phase relative to the c phase, so the temperature at which the e phase forms increases with increasing Si content. In addition, two IMC phases (the r and Laves phases) appear under equilibrium condi- tions over wide temperature and composition ranges. The r phase coexists with the c phase at low Si contents, and the c/c + r boundary gradually decreases to 1320 K as the Si content increases up to 1 wt%. The Laves phase becomes more stable at higher Si contents. In the present study, Si was added in order to replace the r phase with the Laves phase, which has a less complex crystal structure. In addition, Si increases the high-temperature oxidation resistance of Co-based alloys [23]. Therefore, a sufficient amount of Si should be incorporated when thermomechanical processes will be performed, and here, a Si content of 1 wt% was added to extend the temperature range of the single-c-phase region as low as possible. Then, the dependence of the constituent phases on the carbon content in the Co–28Cr–9W–1Si–C alloy was evaluated (Fig. 1b). Carbon doping has been commonly employed to precipitate hard carbide particles in Co–Cr–Mo-based orthopedic alloys [24–26]. Like Si addition, carbon addition also suppresses the formation of the r phase and consequently causes precipitation of the Laves a b Fig. 1. Vertical sections of phase diagrams calculated using Thermo-Calc software for (a) Co–28Cr–9W–Si and (b) Co–28Cr–9W–1Si–C (wt%) systems. 988 K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 phase. Furthermore, carbon forms M 23 C 6 -type carbides, as in the Co–Cr–Mo–C system. Therefore, based on the above consider- ations, this study focused on Co–28Cr–9W–1Si alloys with C con- tents of less than 0.10 wt%, which have a single c phase at a hot working temperature of about 1473 K. 3. Materials and methods 3.1. Specimen preparation Co–28Cr–9W–1Si (wt%) alloys with 0.03 and 0.06 wt% C were prepared in a high-frequency induction furnace in an argon atmo- sphere. Table 1 lists the chemical compositions of the two alloys. Hereafter, the alloys with 0.03 and 0.06 wt% C are referred to as the 0.03C and 0.06C alloys, respectively. Cast ingots (15 mm in diameter and 200 mm long) were subjected to homogenizing heat treatment at 1473 K for 21.6 ks (6 h) and were then processed by multipass hot caliber rolling (U15 mm ? U9.6 mm), followed by water quenching. The equivalent strain, e eq , in the rolling process was calculated as 0.89 using the equation: e eq ¼ ln A 0 A ð1Þ Here, A 0 and A are the cross-sectional areas before and after hot rolling, respectively. 3.2. Microstructural characterization The constituent phases were identified by X-ray diffraction (XRD; Panalytical, X’Pert MPD) using Cu Ka radiation. Solute segre- gation and precipitation were examined using a field-emission scanning electron microscope (FE-SEM; Carl Zeiss, Ultra 55) with an angle-selective backscattered electron (AsB) detector and a field-emission electron probe microanalyzer (FE-EPMA; JEOL, JXA-8430F). The FE-SEM and FE-EPMA were operated at an acceler- ation voltage of 15 kV. Electron backscatter diffraction (EBSD) scans were performed on a FE-SEM (FEI, XL30S-FEG) operated at 20 kV. The EBSD data were accumulated and analyzed using an ori- entation-image microscopy system (TexSEM Laboratories, Inc.). A step size of 0.1 lm was employed in a hexagonal scan grid. In the EBSD analysis, measured points with confidence indices less than 0.1 were eliminated to reduce the inaccuracy in the EBSD measurements and analysis. These points are depicted in black in the constructed EBSD maps. The average grain sizes were mea- sured from the EBSD maps using the line intercept method with a correction factor of 1.128. Samples for XRD, SEM, EPMA, and EBSD were prepared by mechanical grinding and polishing using emery papers and a 0.3-lm alumina suspension, followed by mir- ror polishing with a 0.04-lm colloidal silica solution. TEM observa- tions were performed on a Topcon EM002B operated at 200 kV. Energy-dispersive X-ray spectroscopy (EDS) was performed to determine the chemical compositions of precipitates and the c ma- trix. TEM specimens were produced by cutting 3-mm-diameter disks from each specimen and grinding them to form thin foils using a dimple grinder (Gatan, Model 656). Finally, thin foils were prepared by ion-beam milling (Gatan, Model 691, PIPS). 3.3. Tensile tests Uniaxial tensile tests were performed at room temperature. Tensile specimens (gauge section 1.6 Â 1.0 mm 2 ; gauge length 10.5 mm) were prepared by electrical discharge machining with their longitudinal axis parallel to the rolling direction (RD). Since the ingots used in this preliminary work were relatively small, the tensile specimens did not satisfy any ASTM/ISO standards. Specimens were strained to failure at a nominal strain rate of 1.6 Â 10 À4 s À1 . Tensile tests were performed at least three times for each specimen, and averages and standard deviations (SD) of the 0.2% proof stress, the ultimate tensile strength, and the elonga- tion-to-failure were calculated. The fracture surfaces of the alloys were observed by FE-SEM (Carl Zeiss, Ultra 55). 4. Results The 0.03C and 0.06C alloys could be subjected to high-temper- ature deformation without cracking their surfaces or their interi- ors. This implies that the e phase did not form during hot rolling of the alloys, as intended. The microstructural evolution and its ef- fect on the mechanical properties are described below. Table 1 Chemical compositions of alloys used in the present study (wt%). Alloy Co Cr W Si Ni Mn O N C 0.03C Bal. 27.7 9.06 1.01 0.005 <0.001 0.052 0.028 0.031 0.06C Bal. 27.5 9.02 1.02 0.006 <0.001 0.029 0.025 0.059 a b Fig. 2. XRD spectra of Co–28Cr–9W–1Si–C alloys for three different states, indicating the effect of solution heat treatment and hot rolling on e martensite formation during quenching for (a) 0.03C and (b) 0.06C alloys. K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 989 4.1. Constituent phases Fig. 2a and b respectively shows XRD spectra of the 0.03C and 0.06C alloys in three states (i.e., as-cast, homogenizing heat-trea- ted, and hot-rolled specimens). These XRD results indicate that, de- spite their different carbon contents, similar phases formed in both alloys under each condition. The as-cast specimens exhibit only diffraction peaks assigned to the c phase. In contrast, both c and e reflections are observed for the heat-treated specimens. Thus, the e phase must form as a result of an athermal martensitic trans- formation that occurs during cooling after the heat treatment. Only c reflections are observed in the hot-rolled specimens of both al- loys. The c phase of the hot-rolled alloys is thus metastable (i.e., mechanically stabilized) and would transform to the e phase through the strain-induced martensitic transformation at temper- atures for which the e phase is more stable than the c phase (below 1150 K, Fig. 1). Similar mechanical stabilization [27] of the c phase has been reported in Co–Cr–Mo alloys [13,15,16]. Peaks due to other phases besides the c and e phases are hardly visible in the XRD patterns of any specimen. 4.2. Microstructures and elemental distributions To clarify the effects of the thermomechanical treatment, including both homogenizing heat treatment and subsequent hot rolling, the grain structures and solute segregation of the initial as-cast and hot-rolled alloys were compared. 4.2.1. As-cast specimens Fig. 3a, b and c, d respectively shows EBSD maps of as-cast 0.03C and 0.06C alloys: Fig. 3a and c shows inverse pole figures (IPFs) and Fig. 3b and d shows boundary maps. The black, green, and red lines in the boundary maps indicate high-angle boundaries (HABs) with misorientations larger than 15°, low-angle boundaries (LABs) with misorientations of 2–15°, and annealing twin boundaries (ATBs) with a R3 coincidence site lattice relationship, respectively. Both alloys exhibit cellular dendritic microstructures. Carbon microal- loying was found to significantly reduce the grain sizes of the solid- ification microstructures. 1 In particular, the grain sizes were 81.6 ± 8.3 lm and 45.6 ± 3.6 lm for the 0.03C and 0.06C alloys, respectively. Almost all the grains are surrounded by HABs, but some LABs form because of impingement of crystals with similar orienta- tions. Few ATBs can be identified in the as-cast specimens, although a c b d Fig. 3. EBSD maps of (a and b) 0.03C and (c and d) 0.06C alloys in the as-cast state: (a and c) IPF and (b and d) boundary maps. 1 Although the grain sizes of the as-cast specimens depend on the solidification conditions (e.g., undercooling), the same casting conditions were used throughout this study. Consequently, this grain refinement must originate from carbon doping. The effect of the carbon content on the refinement of the solidification microstruc- tures will be discussed elsewhere. 990 K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 they develop in the c grains in the hot-rolled specimens, as shown below. Fig. 4a and b shows SEM–backscattered electron (BSE) images of the as-cast 0.03C and 0.06C alloys. The grain structures were visu- alized by electron channeling contrast imaging. A high contrast can be observed within the grains. Since BSE imaging discriminates be- tween microstructures based on the mean atomic number, the bright regions observed inside the grains correspond to tungsten segregation. Furthermore, W-rich precipitation is identified in the regions with high W concentrations (Fig. 4c). The 0.06C alloy con- tains more precipitates than the 0.03C alloy. Fig. 5 shows SEM images and EPMA elemental maps of the as- cast specimens, which reveal that Cr and Si are also segregated in regions with high W concentrations. Carbon microalloying en- hances the inhomogeneity of the distributions of such solute ele- ments and increases the precipitate density. The segregation and resulting precipitation would refine the solidification microstruc- tures, as shown in Fig. 3. Pronounced carbon segregation or large quantitates of the carbide M 23 C 6 were not identified. A tiny quan- tity of oxide inclusions (e.g., SiO 2 ) was also identified. Such parti- cles may be formed during melting, because Si has a high affinity for oxygen. 4.2.2. Hot-rolled specimens Fig. 6a, d and b, e respectively shows IPF and boundary maps of cross-sectional areas of the hot-rolled 0.03C (Fig. 6a and b) and 0.06C (Fig. 6d and e) alloys. These specimens have quite similar microstructures and textures, despite their different carbon con- tents. Equiaxed grains surrounded by HABs are frequently ob- served. There are ATBs inside the grains, some of which deviate far from the ideal R3 relationship, resulting in further grain refine- ment. These deviations are caused by interactions between the twin boundaries and lattice dislocations [17]. The grain sizes calcu- lated by neglecting ATBs (i.e., considering only the black lines in the boundary maps) for the 0.03C and 0.06C alloys are 11.1 ± 1.7 lm and 11.8 ± 2.0 lm, respectively. Thus, it is clear that hot rolling produces finer grains than those in the as-cast speci- mens. The local strain evolution in c grains was evaluated using the kernel average misorientation (KAM) approach. The KAM value represents an average misorientation angle for all adjacent mea- surement points in a grain, and it is correlated with the density of geometrically necessary dislocations [28]. Fig. 6c and f shows the results for the 0.03C and 0.06C alloys, respectively. They reveal that the hot-rolled microstructures are distorted, although there are some c grains with low KAM values. The TEM bright-field (BF) images in Fig. 7a and b shows the c grains of the 0.03C and 0.06C alloys, respectively. Selected area dif- fraction (SAD) patterns of the fcc structure obtained with a beam incidence of ½ 110 c are superimposed on each figure. There are no significant differences between the TEM images of the 0.03C and 0.06C alloys. Both exhibit many fringes originating from dislo- cations that are dissociated into stacking faults (SFs) bounded by Shockley partial dislocations. These extended dislocations must strengthen the alloys. Fig. 8a shows an SEM-BSE image of the hot-rolled 0.03C alloy, which has a low precipitate density. In contrast, there are more precipitates in the 0.06C alloy (Fig. 8b). Thus, in terms of microstructural evolution, the two alloys with different carbon contents differ only in the amount of precipitates they contain. 50 µm a 10 µm c 50 µm b Fig. 4. SEM-BSE images of as-cast Co–28Cr–9W–1Si–C alloys: (a) 0.03C and (b) 0.06C alloys. (c) Magnified image of interdendritic region indicated in (b). Co Cr C Si W 50 µm SEM a b 50 µm Low High Co Cr C Si W SEM Fig. 5. SEM images and corresponding EPMA elemental maps of as-cast Co–28Cr–9W–1Si–C alloys: (a) 0.03C and (b) 0.06C alloys. K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 991 TEM analysis was conducted to investigate the crystal struc- tures of the precipitates in the hot-rolled specimens. Fig. 9a shows a BF image of the hot-rolled 0.06C alloy, which includes a precipi- tate. The particle diameters are a few hundred nanometers. The dark-field (DF) image (Fig. 9b) was obtained using the reflection from the precipitate (indicated by the arrow in the SAD pattern in Fig. 9c). The SAD pattern and its key diagram (Fig. 9d) reveal that the precipitates in the hot-rolled specimens have the hexagonal C14 Laves phase. In the DF image (Fig. 9b), lattice defects are apparent in the interior of the Laves phase. In addition, many SFs decollate the precipitate, as shown in the BF image (Fig. 9a). Fig. 10 shows a SEM-BSE image and the corresponding EPMA maps of the hot-rolled 0.06C alloy. These results indicate that the solute segregation induced during casting disappears during the a b d c e f 50 µm 50 µm 50 µm 50 µm 50 µm 50 µm 001 111 101 Fig. 6. EBSD maps of (a–c) 0.03C and (d–f) 0.06C alloys processed by hot rolling: (a and d) IPF, (b and e) boundary, and (c and f) KAM maps. b a Fig. 7. TEM BF images of (a) 0.03C and (b) 0.06C alloys processed by hot rolling. Corresponding SAD patterns of the matrix phase (beam incidence//½ 110 c ) are superimposed in each figure. 992 K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 homogenizing heat treatment and subsequent hot rolling. Further- more, the Laves phase is enriched with W, so it is brighter than the r phase in the SEM-BSE image. Micron-sized, Cr-rich r phase par- ticles were also observed, but their volume fraction was very small. There was no carbide precipitation for the alloys with either carbon content. 4.3. Tensile properties Fig. 11a and b shows nominal stress–nominal strain curves ob- tained from tensile tests of each alloy in the three different states. All the stress–strain curves show uniform elongation followed by sudden fracture without macroscopic necking. This type of fracture behavior is typically observed in Co–Cr–Mo alloys [14,15,17]. The tensile properties obtained by analyzing the stress–strain curves are summarized in Fig. 12 and Table 2. Compared to that of the as-cast specimens, the strength is reduced by homogenizing heat treatment and is drastically improved by hot rolling. Interestingly, hot rolling also enhances the elongation-to-failure. It should be noted that the tensile properties of the as-cast specimens prepared using the current casting technique and the heat-treated specimens were very poor and do not meet the ISO requirements. Table 2 also lists the tensile properties of Co–Cr–Mo alloys subjected to thermo- mechanical processing [14,15,17,21]. The tensile ductility of hot- rolled 0.06C alloy (46 ± 2.4%) is much more remarkable than those of Co–29Cr–6Mo alloys (<30%). Thus, the present approach can be used to obtain novel Co–Cr–W-based alloys with excellent combi- nations of strength and ductility. Furthermore, carbon addition in- creases both the strength and the elongation. 5. Discussion 5.1. Microstructure evolution during thermomechanical processing The present study revealed that hot deformation causes grain refinement and generates a high density of lattice defects. In a par- allel study, the influence of the deformation conditions on the microstructure evolution during hot compression of the developed alloy was systematically investigated [29]. It was revealed that DRX occurs during hot deformation, and this is considered a key metallurgical process for simultaneously realizing dislocation hardening and grain refinement. It should be noted that controlling the stacking fault energy (SFE) is very important for understanding the DRX behavior. In particular, a lower SFE is desirable for achiev- ing the fine grains and homogeneous microstructures after DRX [16,30]. Increasing the Cr concentration reduces the SFE of the al- loys, whereas the addition of the c-stabilizing elements such as W and Ni increases the SFE. Actually, an Co–20Cr–15W–10Ni (L-605) alloy with a higher SFE than the present alloys does not exhibit a high fraction of fine grains (<10 lm) [30]. It is clear that further grain refinement in the present Co–28Cr–9W-based alloys can be realized by optimizing the hot deformation conditions (temperature, strain rate, and strain path), since the smallest grain size obtained in the hot compression test was 1.4 ± 0.3 lm [29]. Therefore, the alloy compositions studied here, namely, high Cr contents and low amounts of c-phase stabilizers (Ni and W), are desirable for grain refinement. In addition, static recrystallization (SRX) was observed to occur during cooling after deformation in the previous study [29]. Similarly, the grains with low KAM values in Fig. 6 indicate the occurrence of SRX during cooling and/or reheating between rolling passes. Thus, the hot-rolled microstructures exhibit local strains that vary from grain to grain. Nevertheless, it should be noted that the mixed DRX/SRX microstructures still remain uniform and fine. 5.2. Precipitation In our alloy design, Si was selected as a microalloying element to cause the brittle r phase to be replaced with the Laves phase. It should be noted that lattice defects were observed inside Laves phase particles (Fig. 9b), which means that in contrast to the brittle r phase, the Laves phase may plastically deform along with the surrounding c phase matrix during hot deformation, although the intermetallic Laves phase is generally hard and brittle. In this situation, the frequency of crack nucleation at the c phase/precip- itate interface will be reduced. Thus, Si addition seems to improve the hot workability. The findings of the present study regarding the precipitation behavior have important implications for carbon microalloying. According to the calculated phase diagram (Fig. 1b), carbon also inhibits the formation of the r phase and leads to precipitation of the Laves phase. The experimental results confirmed this predic- tion and revealed that the amount of Laves particles increases with increasing carbon content. It was reported that the transition from the intermetallic r phase to M 23 C 6 carbide by adding carbon also occurs according to the global reaction: r + C ?M 23 C 6 [31]. How- ever, since the carbide was not identified in this work, the suppres- sion of the r phase formation can be explained in terms of thermodynamics. W-rich precipitates were also observed in the as-cast speci- mens, and their fraction increased with carbon concentration (Fig. 4). Therefore, the Laves phase precipitates in the hot-rolled specimens may be related to solidification segregation. To investi- gate the nature of the Laves phase, the chemical composition of the a b Fig. 8. SEM-BES images of hot-rolled Co–28Cr–9W–1Si–C alloys: (a) 0.03C and (b) 0.06C alloys. K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 993 a c b d Fig. 9. TEM analysis of precipitates in hot-rolled 0.06C alloys: (a) BF image, (b) SAD pattern, (c) DF image obtained using diffraction in (b), and (d) key diagram of (b). Fig. 10. SEM-BSE image and corresponding EPMA elemental maps of hot-rolled Co–28Cr–9W–1Si–0.06C alloy. 994 K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 precipitates and the c matrix in the hot-rolled 0.06C alloy were investigated by performing TEM-EDS analysis (Table 3). Compared with the surrounding c matrix, the Laves phase has higher W and Si contents and lower Co and Cr contents. This is inconsistent with the solute partitioning tendency in the as-cast alloys: the interden- dritic regions where precipitates nucleate have higher Cr contents, in addition to the higher W and Si contents, as compared to the c matrix (Fig. 5). In other words, the precipitates in the as-cast and hot-rolled specimens have different chemistries, and redistribution of the alloying elements occurred during the thermomechanical processing. In contrast, the r phase has higher concentrations of Cr, W, and Si than the c phase (Table 3). Thus, solidification segre- gation is considered to be the origin of the r phase observed in the hot-rolled alloys. n i a r t s l a n i m o N n i a r t s l a n i m o N N o m i n a l s t r e s s ( M P a ) hot-rolled as-cast heat-treated hot-rolled as-cast heat-treated a b Fig. 11. Tensile stress–strain curves of Co–28Cr–9W–1Si–C alloys for three different states (as-cast, heat-treated, and hot-rolled): (a) 0.03C and (b) 0.06C alloys. 0.03C 0.06C a b c 0.03C 0.06C 0.03C 0.06C Fig. 12. Changes in (a) 0.2% proof stress, (b) ultimate tensile strength, and (c) elongation-to-failure of the Co–28Cr–9W–1Si–C alloys. The 0.2% proof stress, ultimate tensile strength, and elongation-to-failure were obtained from the nominal stress–nominal strain curves in Fig. 11. Table 2 Tensile properties of Ni-free Co–28Cr–9W–1Si–C alloys obtained after multipass hot rolling. The results for Ni-free Co–29Cr–6Mo alloys subjected to thermomechanical processing and Co–20Cr–15W–10Ni (L-605) alloy are also shown for comparison. Note that the initial Co–29Cr–6Mo alloy specimens used in Refs. [14,15] are not fully annealed and therefore have higher strength than the annealed Co–28Cr–9W-based alloys. Alloys Condition / method Equivalent strain imposed 0.2% proof stress (MPa) Ultimate tensile strength (MPa) Elongation (%) Yield ratio References Co–28Cr–9W–1Si–0.03C As-cast – 387 ± 16 854 ± 55 27.5 ± 6.4 0.45 Present work Annealed – 277 ± 23 689 ± 30 31.8 ± 1.3 0.40 Hot-rolled 0.89 500 ± 31 1145 ± 46 35.2 ± 3.1 0.44 Co–28Cr–9W–1Si–0.06C As-cast – 350 ± 10 806 ± 36 27.2 ± 3.9 0.43 Present work Annealed – 328 ± 23 926 ± 8 39.6 ± 1.7 0.35 Hot-rolled 0.89 564 ± 13 1282 ± 12 46.0 ± 2.4 0.44 Co–29Cr–6Mo Initial – 570 980 16.8 0.58 [14] Hot-forged 0.9 800 1320 14.5 0.61 Hot-forged 1.5 1050 1450 10.3 0.72 Hot-forged 1.8 1330 1450 2.5 0.92 Co–29Cr–6Mo Hot-forged 1.1 590 718 15.1 0.82 [21] Hot-forged 1.7 648 1050 22.6 0.62 Hot-forged 3.5 890 1415 21.6 0.63 Co–29Cr–6Mo–0.12N Initial – 540 988 25.2 0.55 [15] Hot-forged 1.8 1400 1620 21.8 0.86 Co–29Cr–6Mo–0.14N Hot-rolled 3.1 1080 1415 20.6 0.76 [17] Co–20Cr–15W–10Ni (L- 605) Annealed – 550 ± 5 1118 ± 5 39 ± 1 0.49 [37] K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 995 5.3. Mechanical characteristics of the developed alloys 5.3.1. Strengthening mechanism As mentioned above, the DRX during the hot rolling causes grain refinement and the accumulation of lattice defects (Figs. 6 and 7). The yield stress increments derived from the current hot rolling process are 113 MPa and 214 MPa for the 0.03C and 0.06C alloys, respectively. However, supposing that the Hall–Petch rela- tionship holds for the Co–Cr–Mo alloys [15], the grain refinement effect only occurs at 60 MPa and is considered not to be domi- nant. In addition, solid solution hardening due to interstitial carbon is not considered to be significant for biomedical Co–Cr-based al- loys [25]. On the other hand, the developed alloys contained submicron- sized Laves-phase particles, since carbon doping increases the amount of Laves phase precipitated in the c matrix of hot-rolled specimens (Fig. 8). The intermetallic Laves phase would be much harder than the surrounding c matrix. To examine its effect, the well-known Orowan mechanism can be used to predict the strength increase due to impenetrable particles according to the following expression [32]: Dr Orowan ¼ aGb ffiffiffi f p d ln d 2b ð2Þ where Dr Orowan is the strength increase, a is a constant, G is the shear modulus, b is the magnitude of the Burgers vector, and f and d are respectively the volume fraction and diameter of the par- ticles. The image analysis revealed that the area fraction of the Laves phase in the hot-rolled 0.06C alloy specimen is less than 1%, and therefore f in Eq. (2) will be very small. Thus, the Laves particles themselves seem to make a negligible contribution to strengthen- ing. Again, the hardness difference between the c phase and the Laves particles will directly increase the dislocation density during hot deformation. Therefore, in summation, dislocation hardening would play the most significant role, as for Co–Cr–Mo alloys [17]. In particular, a possible explanation for the higher yield strength of the 0.06C alloy is the increased dislocation density induced by deformation of the c–Laves duplex microstructures. 5.3.2. Enhanced tensile ductility It is well known that the ductility of metallic materials can be simply explained in terms of the plastic instability, which is ex- pressed by the Considére criterion: r P dr de ð3Þ where r, e, and dr/de are the true stress, true strain, and work hard- ening rate obtained in tensile tests, respectively. Eq. (3) implies that when the work hardening rate decreases to become equal to the flow stress (plastic instability), macroscopic necking followed by facture will occur in the specimens. Fig. 13a shows the true stress and work hardening rate curves of the hot-rolled 0.03C and 0.06C alloys as a function of the true strain. Both alloys have nearly iden- tical work hardening rates during tensile tests, and fractures occur a little before the onset of plastic instability for both 0.03C and 0.06C alloys, which is similar to the behavior of biomedical Co–Cr–Mo al- loys [15,17,33]. Nevertheless, the elongation-to-failure of the 0.06C alloy is significantly larger than that of the 0.03C alloy. This clearly indicates that carbon microalloying is very effective in improving the ductility of the alloys within the carbon concentration range investigated (<0.1 wt%). However, further investigation is required to clarify its mechanism. In order to understand the plastic deformation behavior of the present alloys, the fracture surfaces of the tensile-tested specimens were investigated. There was no significant difference between the alloys. Fig. 13b shows the facet planes of the hot-rolled 0.06C al- loys, some of which have striations (makings). These features indi- cate quasi-cleavage-type and/or intergranular-type fractures, which are related to e martensite [15,34]. Therefore, the plastic deformation mechanism of the present Co–28Cr–9W-based alloys is very similar to that of the Co–Cr–Mo orthopedic alloys: it in- volves planar slip of Shockley partial dislocations and is hence a strain-induced martensitic transformation (SIMT). This is reason- able, because the c phase in the hot-rolled alloy specimens is con- sidered metastable (Fig. 2). Yamanaka et al. reported that the elongation of Ni-free Co–29Cr–6Mo alloys increases with reduc- tions in grain size and reaches a maximum when the grain size equals approximately 5.0–10 lm [14]. Thus, the effect of grain refinement on the ductility enhancement should be taken into ac- count for the current alloys. 5.3.3. Effects of static recrystallization on mechanical characteristics The hot-rolled Co–28Cr–9W–1Si–C alloys exhibit SRX grains with low strains (Figs. 6 and 7). A reduction in hardness due to the introduction of SRX grains would be preferred for dental appli- cations, because the amount of chipping failure occurring in the machining process increases with increasing hardness [35]. The occurrence of SRX must contribute to the tensile properties of the alloys, as well. Dini et al. reported that cold-rolled and subsequently annealed Fe–31Mn–3Al–3Si austenitic steels with partially recrystallized microstructures show excellent combina- Table 3 Chemical compositions of Laves phase, r phase, and c matrix in hot-rolled Co–28Cr– 9W–1Si–0.06C alloy obtained from TEM-EDS analysis. Standard deviations are given in parentheses. Phase Unit Co Cr W Si Laves (wt.%) 28.45 12.68 55.61 3.26 (0.67) (0.47) (1.59) (0.48) (at.%) 42.17 21.29 26.25 10.09 (0.24) (0.30) (1.45) (1.29) Sigma (wt.%) 45.58 37.03 17.35 1.36 (0.27) (0.29) (0.46) (0.03) (at.%) 46.73 43.13 5.71 2.93 (0.28) (0.20) (0.12) (0.08) c matrix (wt.%) 60.85 27.84 10.14 1.16 (0.49) (0.15) (0.04) (0.60) (at.%) 62.04 32.17 3.31 2.48 (0.90) (0.37) (0.01) (1.28) a b Fig. 13. (a) True stress–true strain and work hardening rate–true strain curves of hot-rolled Co–28Cr–9W–1Si–C alloys and (b) fracture surface of hot-rolled 0.06C alloy. 996 K. Yamanaka et al. / Materials and Design 55 (2014) 987–998 tions of strength and ductility [36]. In particular, they found that the unrecrystallized areas increase the strength, while the ductility increases with increasing fractions of the recrystallized areas. Sim- ilarly, the formation of SRX grains observed here lowers the yield stress and simultaneously enhances the ductility. Although the present alloys were designed for dental applications, such materi- als are also considered suitable for balloon-expandable-stent struts. Certain stents undergo a deformation of approximately 40% during their expansion [37]. Thus, the Co–20Cr–15W–10Ni al- loy (known as L-605 or ASTM F90), which has excellent ductility, has been actually used for vascular stents because of its excellent elongation-to-failure, in spite of its high concentration of nickel. Table 2 also shows the typical tensile properties of the L-605 alloy [38], which are comparable to those of the hot-rolled 0.06C alloy. A higher tensile strength and a lower yield ratio (=yield stress/tensile strength) are preferred for stent materials [39]. Furthermore, cur- rent trends in coronary stent design are towards thinner struts, and Co–Cr alloys with high elastic moduli (>200 GPa) are desirable for ultrathin struts [39]. In this case, the grain sizes of the alloys must be reduced to improve the mechanical performance of the stent struts [40]. Therefore, Ni-free Co–28Cr–9W–1Si–C alloys with fine-grained microstructures, low yield ratios, and excellent ductil- ity are promising for dental as well as stent materials. These micro- structural and mechanical characteristics are difficult to obtain with Co–Cr–Mo alloys, because their extremely low SFE restricts the SRX as well as the recovery process [17], resulting in very high yield stress/hardness (Table 2). 6. Conclusions A microalloying and thermomechanical processing strategy for designing Co–Cr–W-based alloys for the disk materials used in CAD/CAM machining has been presented. The alloy composition of the commercialized Co–28Cr–9W alloy used in the as-cast state was modified based on thermodynamic calculations. Then, the thermomechanical processing was performed to induce grain refinement due to DRX and a high density of SFs, and, conse- quently, enhanced strength was obtained. This processing also yields uniform elemental distributions. The microalloying ele- ments Si and C were selected so that the brittle r phase would be replaced with the Laves phase, which is plastically deformable at elevated temperatures. It was found that with increasing carbon content, the precipitation of IMC Laves-phase particles in the hot- rolled specimens increased. The fine precipitates in the c matrix were found to negligibly affect the strength through precipitation hardening. However, the Laves phase particles may accelerate the accumulation of lattice defects during hot deformation, resulting in strengthening of the hot-rolled alloys. This approach increased both the alloy strength and the ductility, and the resulting novel Co–Cr–W-based alloy processed by hot rolling is promising as a disk material for use with dental CAD/CAM technology as well as for stent struts. Acknowledgements The authors would like to thank Isamu Yoshii, Kimio Wako, Fumiya Sato, Shun Ito, and Issei Narita, IMR, Tohoku University, for contributing sample preparation, TEM observations and EPMA analysis. This research was financially supported by the Grant- in-Aid for JSPS Fellows, the Global COE Program ‘‘Materials Integra- tion (International Center of Education and Research), Tohoku University’’, the Regional Innovation Strategy Support Program, NICHe, Tohoku University, and the Regional Innovation Cluster Program from the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. References [1] Takaichi A, Suyalatu, Nakamoto T, Joko N, Nomura N, Tsutsumi Y, et al. 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