journal of the mechanical behavior of biomedical materials 29 (2014) 417–426
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Research Paper
Effects of nitrogen addition on microstructure and mechanical behavior of biomedical Co–Cr–Mo alloys Kenta Yamanakaa,n, Manami Morib, Akihiko Chibaa a
Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan NISSAN ARC, LTD., Yokosuka 237-0061, Japan
b
art i cle i nfo
ab st rac t
Article history:
In the present study, the microstructures and tensile deformation behaviors of biomedical
Received 19 July 2013
Co–29Cr–6Mo (wt%) alloys containing different concentrations of nitrogen (0–0.24 wt%)
Received in revised form
were systematically investigated. As the nitrogen concentration increased, the volume
28 September 2013
fraction of athermal ε martensite decreased, because nanoprecipitates hindered the
Accepted 7 October 2013
formation of stacking faults (SFs) by acting as obstacles to Shockley partial dislocation
Available online 23 October 2013
formation, and athermal ε martensite usually forms through the regular overlapping of
Keywords:
SFs. The formation of the athermal ε martensite was completely suppressed when the
Biomedical Co–Cr–Mo alloys
nitrogen concentration exceeded 0.10 wt%, resulting in a simultaneous improvement in
Nitrogen addition
the strength and ductility of the alloys. It was found that the glide of the Shockley partial
Microstructures γ-ε martensitic transformation Mechanical properties
dislocations and the strain-induced γ (fcc)-ε (hcp) martensitic transformation (SIMT) operated as the primary deformation mechanisms. However, adding nitrogen reduced the work hardening by suppressing the formation of the SFs and preventing the SIMT from taking place. This resulted in an intrinsic decrease in the tensile ductility of the alloys. It is also shown that all the alloys exhibited premature fractures owing to the SIMT. The formation of annealing twins in the γ grains is found to be enhanced by nitrogen addition and to promote the SIMT, resulting in a reduction in the elongation-to-failure due to nitrogen addition. These results should aid in the design of alloys that contain nitrogen. & 2013 Elsevier Ltd. All rights reserved.
1.
Introduction
Co–Cr–Mo alloys exhibit excellent biocompatibility, corrosion resistance, and wear resistance, and are therefore widely used for orthopedic implants such as artificial hip and knee joints (Niinomi, 2002; Buford and Goswami, 2004; Chiba et al., 2007). It is known that metal-on-metal bearings made of Co–Cr–Mo alloys allow large-diameter femoral head components to be used, which provide a greater range of motion than that possible with conventional artificial hip joints n
Corresponding author. Tel.: þ81 22 215 2118; fax: þ81 22 215 2116. E-mail address: [email protected] (K. Yamanaka).
1751-6161/$ - see front matter & 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jmbbm.2013.10.006
(Cuckler et al., 2004). Thus, further improvements in the characteristics of these alloys, especially in their wear resistance and mechanical durability, are essential to expand further the biomedical applications of these alloys, and it is of great importance to understand the mechanisms of their functions. In Co–Cr–Mo alloys that comply with the ASTM F75 standard, up to 0.35 wt% carbon can be incorporated to allow for carbide precipitation, which is a major strengthen ing mechanism (Rajan, 1982; Caudillo et al., 2002; Mineta et al., 2010). However, recent studies have revealed that hard
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journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
carbide precipitates are sometimes detrimental to the wear resistance, corrosion resistance, and biocompatibility of the alloys (Chiba et al., 2007; Battini et al., 2011; Liao et al., 2012). Therefore, significant efforts are being made to formu late novel alloys and devise hot deformation processes for these alloys (Salinas-Rodriguez and Rodriguez-Galicia, 1996; Escobedo et al., 1996; Kilner et al., 1987; Hsu and Lian, 2003; Chiba et al., 2007; Matsumoto et al., 2010; Mani, SalinasRodriguez, López, 2011; Yamanaka et al., 2009, 2011, 2012a, 2012b; Mori et al., 2010, 2012; Lee et al., 2008a). The matrix phase also significantly influences the mechanical and tribological properties of Co–Cr–Mo alloys. Generally, these alloys consist of a metastable γ (fcc) matrix and/or platelike ε (hcp) martensite. The ε martensite is formed during quenching, plastic deformation, and isothermal heat treatments and is known to enhance the strength and wear resistance of the alloy (Chiba et al., 2007; Mani et al., 2011). However, the formation of ε martensite also leads to poor deformability (Mori et al., 2010; Mani et al., 2011; Yamanaka et al., 2011, 2012a). Nickel addition is usually employed to prevent such fractures due to the presence of the ε martensite in cobalt-based alloys. In fact, biomedical Co–Cr alloys containing large amounts of Ni (10–37 wt%) (conforming to ASTM F90 and F562) exhibit excellent ductility and have been used in a number of practical applications (Nagai et al., 2012; Marrey et al., 2006). However, nickel is known to cause allergies and cancer in living organisms (Denkhaus and Salnikow, 2002). Nitrogen is known to be an alloying element that can replace nickel. It also stabilizes the γ phase and improves the mechanical properties of Co–Cr–Mo alloys. Early studies have reported that the addition of nitrogen to Co–Cr–Mo alloys results in improvements in the tensile strength, plastic deformability, fatigue strength, and wear properties of the alloys (Escobedo et al., 1996; Hsu and Lian, 2003). Recently, Lee et al. (2008a) and Yoda et al. (2012) investigated the microstructures and tensile properties of as-cast Co–Cr–Mo– N alloys and found that alloys with high nitrogen contents exhibit good combinations of strength and ductility that make them suitable for dental applications. However, the studies performed so far either employed specimens containing relatively large amounts of nickel ( 2.5 wt%) and carbide precipitates (carbon contents 0.45 wt%) (Escobedo et al., 1996; Hsu and Lian, 2003) or varied the chromium concentrations of the samples in order to change their nitrogen concentrations (Lee et al., 2008a; Yoda et al., 2012), which prevents the understanding of the intrinsic effects of nitrogen in Co–Cr–Mo alloys. On the other hand, we have previously reported that Ni- and C-free Co–29Cr–6Mo–N (wt%) alloys exhibit high tensile ductility as well as cold workability while
maintaining their high strength (Lee et al., 2008a; Mori et al., 2010, 2012; Yamanaka et al., 2012b). Nevertheless, a systematic investigation of the effects of nitrogen on the structures and mechanical properties of the Ni- and C-free Co–Cr–Mo biomedical alloys has not been done yet. An important point to note regarding nitrogen-containing Co–Cr–Mo alloys is their decomposition on the nanoscale in the γ phase (Yamanaka et al., 2013). Nanometer-sized nitride precipitates formed in the γ matrix were found to interact with extended dislocations consisting of Shockley partial dislocations and stacking faults (SFs), ultimately inhibiting the γ-ε martensitic transformation (Yamanaka et al., 2013). In this context, the aim of the present study was to investigate the dependence of the microstructures, phase distributions, stabilities of the γ matrices, dislocation structures, and tensile deformation behaviors of Co–29Cr–6Mo–N alloys on the nitrogen concentration.
2.
Materials and methods
2.1.
Specimen preparation
Co–29Cr–6Mo (wt%) alloys with different compositions were prepared in an argon atmosphere using a high-frequency induction furnace. A Cr2N powder was used as the nitrogen source. The chemical compositions of the alloys are shown in Table 1. All the alloys conformed to the ASTM F75 standard and are denoted as 0N, 0.05N, 0.07N, 0.10N, 0.17N, and 0.24N in accordance with the concentration of nitrogen in each alloy. Cast ingots of these alloys of 15 mm in diameter and 200 mm in length were subjected to a homogenizing heat treatment at 1473 K for 1.8 ks and subsequently processed by multipass hot caliber rolling, resulting in a change in their diameter from 15 mm to 9.6 mm. This was followed by the quenching of the hot-rolled alloy bars in water. The equivalent strain, εeq, generated during the rolling process, which was found to be 0.89, was calculated using the following equation: εeq ¼ ln
A0 A
ð1Þ
Here, A0 and A are the cross-sectional areas of the initial and hot-rolled specimens, respectively. Next, the hot-rolled bars were heat treated in air at 1473 K for 0.6 ks and then quenched in water.
Table 1 – Chemical compositions of alloys used in the present study (wt%). Alloy
Co
Cr
Mo
N
Ni
Mn
Si
O
C
0N 0.05N 0.07N 0.10N 0.17N 0.24N
Bal. Bal. Bal. Bal. Bal. Bal.
28.7 27.3 29.0 29.2 27.8 29.9
6.3 6.1 6.0 5.9 6.0 6.1
0.00 0.05 0.07 0.10 0.17 0.24
o0.01 o0.01 o0.01 o0.01 o0.01 o0.01
o0.01 o0.01 o0.01 o0.01 o0.01 o0.01
o0.1 o0.1 o0.1 o0.1 o0.1 o0.1
0.022 0.021 0.021 0.038 0.025 0.028
0.0016 0.0012 0.0012 0.0018 0.0011 0.0028
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
2.2.
Characterization of the microstructures
values were calculated. The fracture surfaces of the alloys were observed using FESEM (Carl Zeiss ULTRA 55).
The microstructures of the alloys were investigated using scanning electron microscopy (SEM), electron-backscatter diffraction (EBSD) analysis, and conventional and high-resolution transmission electron microscopy (TEM and HRTEM, respectively). Cross sections of the specimens, perpendicular to the rolling direction (RD), were observed. The samples subjected to SEM and EBSD analysis were mechanically ground and polished using emery papers and a 0.3 μm alumina suspension. These processes were followed by polishing to a mirror-like finish using a 0.04 μm colloidal silica solution. SEM observations were conducted using a field-emission SEM (FESEM) (Carl Zeiss ULTRA 55) operated at an acceleration voltage of 15 kV. The EBSD measurements were also performed using an FESEM (FEI XL30S-FEG) operated at 15 kV. The EBSD data were collected and analyzed using an orientation image microscope system (TexSEM Laboratories, Inc.). A step size of 0.1 μm in a hexagonal scan grid was employed. TEM and HRTEM were performed using a Topcon EM002B operated at 200 kV. The specimens for TEM were produced by cutting a disk 3 mm in diameter from each alloy and grinding it using a dimple grinder (Gatan Model 656). Next, thin foils were prepared by ion-beam milling these disks (Gatan Model 691, PIPS). To examine the microstructures developed during plastic deformation, TEM observations of the specimens were also performed after they had been subjected to compression tests at a strain rate of 1.0 10 4 s 1 at room temperature, which resulted in a 5% total strain.
2.3.
Tensile tests
Uniaxial tensile tests were performed on the alloy samples at room temperature. Specimens 1.6 1.0 mm2 in the gauge section and 10.5 mm in gauge length were prepared by electrical discharge machining, with the longitudinal axis of the bars parallel to their RD. The samples were strained until failure at a nominal strain rate of 1.6 10 4 s 1. The tensile tests were performed at least three times on each alloy, and the averages and standard deviations (S.D.) of the 0.2% proof stress, ultimate tensile strength, and elongation-to-failure 0.05N
0.07N
3.
Results
3.1.
Microstructures and constituent phases
Fig. 1 shows the EBSD maps of the Co–29Cr–6Mo–N alloys. All the alloys consisted of equiaxed γ grains with a mean grain size of about 100 μm. The phase maps indicate that the 0N, 0.05N, and 0.07N alloys exhibited duplex microstructures comprising a γ matrix and a plate-like ε phase, whereas the 0.10N, 0.17N, and 0.24N alloys exhibited only the γ phase. The ε phase must form through the athermal martensitic transformation occurring during cooling after the heat treatment, because TEM analysis confirmed that both phases have the Shoji–Nishiyama orientation relationship, i.e., (111)γ//(0001)ε and [110]γ//[1120]ε. The inverse pole figure (IPF) and phase maps revealed the formation of multivariant ε martensite structures inside a γ grain, and the number of ε variants decreased with increasing nitrogen content. In the grain boundary (GB) maps, the black, green, and red lines indicate the high-angle boundaries (HABs) with misorientation angles greater than 151, low-angle boundaries (LABs) with misorientation angles of 2–151, and annealing twin boundaries (ATBs) with a Σ3 coincidence site lattice relationship, respectively. The intragranular boundaries in the 0N, 0.05N, and 0.07N alloys were attributable to athermal ε martensite, whereas those in the alloys with over 0.10 wt% nitrogen were due to annealing twins. On the basis of the results of the EBSD analysis, we determined the dependence of the fractions of athermal ε martensite and of ATBs on the nitrogen concentration (Fig. 2). The volume of ε martensite decreased with increasing nitrogen concentration and dropped sharply to zero when the nitrogen concentration was higher than 0.10 wt%. Thus, the critical nitrogen concentration for obtaining a matrix comprising only the γ phase was around 0.10 wt% for the Co–29Cr–6Mo–N system. In contrast, the duplex alloys had 0.10N
0.17N
0.24N
GB
Phase
IPF
0N
100 µm 111 1010 001
101 0001
419
2110
Fig. 1 – EBSD maps of Co–29Cr–6Mo alloys with different nitrogen concentrations.
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Fig. 2 – Volume fractions of athermal ε martensite and ATBs in the Co–29Cr–6Mo alloys as a function of nitrogen concentration.
almost no annealing twins. The volume fraction of ATBs increased rapidly from 0% to 60% for nitrogen concentrations ranging from 0.07 to 0.10 wt%, respectively. In addition, the volume fraction of annealing twins increased gradually from 60% to 75% for nitrogen concentrations ranging from 0.10 to 0.24 wt%, respectively. This was because the formation of multivariant twins within the grains became more pronounced over this range of nitrogen concentrations. Fig. 3a and b shows the SEM–BSE (backscatter electron) images of the 0N and 0.17N alloys with the duplex and single γ matrices, respectively. The microstructure was observed to be subdivided by the athermal ε martensite in Fig. 3a, as well. The s phase, which would appear brighter in these images, was hardly identified in any specimen. Instead, a small number of submicron- and/or micron-sized particles lower in brightness were noted (indicated by the arrow in Fig. 3, for example). The energy dispersive X-ray spectroscopy (EDX) results obtained using the FESEM revealed that these particles were oxides of Cr and Al and might have originated from the raw materials or the alumina crucible used. These oxide particles do not originate from nitrogen addition. They are considered to have a negligible effect on quasi-static tensile deformation behavior discussed below, because their volume fraction is very low.
3.2.
Fig. 3 – SEM–BSE images of (a) 0N and (b) 0.17N alloys subjected to annealing at 1473 K for 0.6 ks, followed by water quenching.
Tensile properties
Fig. 4 shows the nominal stress–nominal strain curves of the Co–29Cr–6Mo–N alloys with nitrogen concentrations ranging up to 0.24 wt%, as obtained from tensile tests. All the specimens were tested to fracture, and macroscopic necking was not noted in any of them. Fig. 5a shows the 0.2% proof stress and ultimate tensile strength of the alloys as a function of their nitrogen concentration (also summarized in Table 2). The 0N alloy exhibited a higher strength than the 0.05N alloy. This may be due to the athermal ε martensite plates, on which the lattice dislocations introduced during water quenching impinge (see Fig. 6a). In contrast, the 0.2% proof
Fig. 4 – Tensile stress–strain curves of Co–29Cr–6Mo alloys with various nitrogen concentrations.
stress was linearly proportional to the nitrogen concentration for concentrations greater 0.05 wt%. It can be seen from Fig. 5a that there was an increase of 1264 MPa in the 0.2% proof stress of the alloys after the incorporation of 1 wt% nitrogen. A similar linear dependence of the strength on the nitrogen concentration has been reported for austenitic steels
421
Elongation (%)
Strength (MPa)
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
Nitrogen content (wt.%)
Nitrogen content (wt.%)
Fig. 5 – Changes in (a) the 0.2% proof stresses and the ultimate tensile strength and (b) uniform elongation as a function of the nitrogen concentration of the Co–29Cr–6Mo–N alloys. The values of the 0.2% proof stress, ultimate tensile strength, and uniform elongation were obtained from the nominal stress–nominal strain curves (Fig. 4). Table 2 – Tensile properties of Co–29Cr–6Mo alloys with different nitrogen concentrations.
0N 0.05N 0.07N 0.10N 0.17N 0.24N
0.2% proof stress (MPa)
Ultimate tensile strength (MPa)
Uniform elongation (%)
Work hardening exponent, n
425758 347724 38176 403736 507715 587734
928747 72874 847722 863741 931711 96279
11.571.9 9.970.6 17.170.3 24.170.5 22.571.2 20.471.1
0.417 0.390 0.384 0.467 0.454 0.405
Fig. 6 – Bright-field TEM images of the Co–29Cr–6Mo alloys with different nitrogen concentrations: (a) 0N and (b) 0.10N alloys.
(Simmons, 1997; Soussan et al., 1991; Lee et al., 2008b). However, the strengthening effect of nitrogen in the present alloys (1264 MPa per wt%) was much more prominent than that in austenitic steels (200–500 MPa per wt%). The ultimate tensile strength decreased and then increased linearly with increasing nitrogen concentration. Fig. 5b shows the dependence of the elongation-to-failure of the Co–29Cr–6Mo–N alloys on the nitrogen content. The 0N alloy exhibited a greater elongation than the 0.05N alloy. On the other hand, it is apparent that eliminating the athermal ε martensite phase (as in alloys with nitrogen contents higher than 0.10 wt%) is an extremely effective way of improving the ductility of the alloys. Since the tensile elongation of the alloys with nitrogen concentrations greater than 0.10 wt%
shows a slight decrease with increasing nitrogen concentration, we can clearly identify the transition area where the elongation-to-failure dramatically changes. The 0.07N alloy falls in this area.
4.
Discussion
4.1. Effect of nitrogen addition on the formation of stacking faults and the athermal γ-ε martensitic transformation Fig. 6a and b shows bright-field TEM images of the 0N and 0.10N alloys, which do and do not contain athermal ε
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Fig. 7 – Bright-field TEM images of the 0.24N alloy annealed at 1473 K for 0.6 ks: (a) before and (b) after compression by 5% at room temperature. HRTEM image taken from the annealed 0.24N alloy specimen is shown in (c).
Table 3 – Equilibrium nitrogen content in the matrix (γ or ε) and Cr2N phases in Co–29Cr–6Mo–(0.1, 0.2)N alloy at various temperatures (wt%) obtained using Thermo-Calc software. Phase
Temperature (K)
Total N content (wt%)
N content in γ or ε phase (wt%)
N content in Cr2N phase (wt%)
Mole fraction of Cr2N phase (%)
γ phase
1473 1473 1273 1273
0.1 0.2 0.1 0.2
8.51 10 3 8.45 10 3 1.10 10 3 1.10 10 3
10.4 10.5 11.0 10.9
1.20 2.49 1.26 2.53
ε phase
1073 1073 873 873
0.1 0.2 0.1 0.2
8.25 10 3 8.25 10 3 0.67 10 3 0.67 10 3
11.3 11.3 11.7 11.7
1.17 2.42 1.24 2.48
martensite, respectively. Dislocations in both alloys were found to have frequently dissociated into SFs, forming highly planar dislocation arrays. Compared with the nitrogenbearing austenitic steels (Ojima et al., 2009), both alloys have very high widths of their SFs. These dislocation structures indicate that the stacking fault energies (SFEs) of the alloys are very low, as predicted by thermodynamic calculations; the SFE values eventually become negative at temperatures lower than 1100 K (Yamanaka et al., 2009, 2012a). Although the SFEs are usually positive for conventional metals and alloys, a negative SFE simply means that the ε phase is more stable than the γ phase. It is well known that SFs can be regarded as local hcp structures embedded in the fcc matrix. Thus, the formation of thermodynamically stable SFs and ε martensite is preferred in the Co–29Cr–6Mo alloys. A similar dislocation structure was also observed in the alloy containing 0.20 wt% N in our previous study (Yamanaka et al., 2013). The addition of nitrogen in greater amounts resulted in significantly different states of the dislocations than those present in the alloys with lower nitrogen concentrations. Fig. 7a shows bright-field TEM images of the 0.24N alloys. Although the black contrast in this figure corresponds to the pairs of Shockley partial dislocations bounding SFs, the width of SFs is not as large as in Fig. 6. In contrast, Fig. 7b shows a bright-field TEM image of the dislocation structures in the 0.24N alloy deformed by 5% during compression at room temperature. It should be noted that the dissociation of the dislocations increased after the deformation and resulted in
the formation of a planar dislocation structure consisting of widely extended SFs. This structure is similar to those of the alloys with lower nitrogen concentrations (Fig. 6). The transition of dislocation structures observed here should be understood in light of nanoprecipitate formation in the γ matrix originating from the nitrogen doping. An HRTEM image of the 0.24N alloy is shown in Fig. 7c. The corresponding selected area diffraction (SAD) pattern for an incident beam parallel to the [110]γ direction is superimposed in the figure. By analyzing the SAD patterns obtained under different incident beam conditions, we identified nanosized plate-like Cr2N precipitates that formed on the {111} planes of the γ phase. Furthermore, streaks originating from the precipitates were identified along the o1114γ directions in the SAD pattern. The nitrogen contents in the γ, ε, and Cr2N phases as a function of temperature and total nitrogen content of the alloys were calculated using the Thermo-Calc software (Table 3). The thermodynamic data sets needed for this calculation were acquired from TCS Steels/Fe-alloys Database Version 6. Note that the ε phase is stable below 1200 K in the Co–29Cr–6Mo–N alloys (Yamanaka et al., 2013). The thermodynamic calculations indicated that the solubilities of nitrogen in the γ and ε phases are both quite low, and therefore the Cr2N phase forms in both the γ and ε phases. In addition, the N concentrations in the Cr2N precipitates are determined by the temperature and are almost independent of the total N content. In other words, increasing the bulk nitrogen content results in an increase in the fraction of Cr2N particles. In addition, the amount of Cr2N phase is
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
423
Fig. 8 – EBSD maps of the microstructures near the fracture surfaces of the (a–d) 0.10N, and (e–h) 0.24N alloys deformed by tensile stress to fracture: (a and e) IPF, (b and f) phase, (c and g) GB, and (d and h) KAM maps.
expected to increase with increasing total N content. The previous study (Yamanaka et al., 2013) revealed that the Cr2N nanoprecipitates act as obstacles for the Shockley partial dislocations and ultimately retard the γ-ε martensitic transformation. Since the amount of the Cr2N phase increases with increasing N concentration (Table 3), the volume fraction of athermal ε martensite consequently decreased with increasing N concentration (Fig. 2). These nanoprecipitates can be overcome under the combined action of the applied stress and thermal energy. Therefore, wide SFs can form even in the 0.24N alloys when a sufficiently high stress is applied (Fig. 7b). Namely, these results indicate that adding nitrogen reduces the planarity in dislocation structures of the Co–Cr–Mo alloys, but the width of SFs becomes larger during tensile deformation and, consequently, the highly planar dislocation structure forms even in the high-nitrogen alloys. Note that perfect dislocations as well as dislocation cells/walls are hardly observed in our alloys investigated in this research, because the SFE of the alloys is negative (Yamanaka et al., 2009, 2012a).
4.2.
Plastic deformation and fracture mechanisms
Fig. 8 shows EBSD maps of cross sections of the tensilefractured specimens of the (a–d) 0.10N and (e–h) 0.24N
alloys. As mentioned above, both alloys contain a negligibly small amount of the athermal ε martensite phase (Fig. 1) and exhibited similar elongations-to-failure (Table 2). In the IPF maps (Fig. 8a and e) and phase maps (Fig. 8b and f) of both alloy specimens, we can identify the in-grain orientation gradient and the ε phase, which is formed via the strain-induced γ-ε martensitic transformation (SIMT) during room-temperature deformation. Thus, for the nitrogen concentrations investigated in this study up to 0.24 wt%, the deformation mode was slipping of the Shockley partial dislocations and subsequent formation of ε martensite, although Yoda et al. (2012) have stated that adding nitrogen changes the deformation mechanism from the SIMT to another form such as twinning or dislocation slipping. Despite these similar IPF and phase maps, the GB maps revealed that quite different microstructures developed during tensile deformation of the two alloys: there were intragranular HABs related to the SIMT in the 0.10N alloys (Fig. 8c), whereas a significant number of deformation-induced LABs developed in the 0.24N alloys (Fig. 8g). Accordingly, the distribution of the local strain was different in the two alloys. The kernel average misorientation (KAM) maps of the alloys are shown in Fig. 8d and h. The KAM value represents an average misorientation angle between all adjacent
424
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
Fig. 9 – SEM images showing typical fracture surfaces of the Co–29Cr–6Mo alloys with different nitrogen concentrations: (a and b) 0N, (c and d) 0.10N, and (e and f) 0.24N alloys.
measurement points within a grain1 and is correlated with the density of geometrically necessary dislocations (Calcagnotto et al., 2010). In the 0.10N alloys (Fig. 8d), the KAM values were preferentially higher in the vicinities of the γ GBs and in the ε martensite plates. These results can be interpreted in terms of GB compatibility and impingement of the dislocations or the plate-like ε phase formed by the SIMT at these boundaries. In contrast, the 0.24N alloys exhibited a more uniform distribution of KAMs inside the γ grains (Fig. 8h). The higher density of Cr2N nanoprecipitates and higher fraction of annealing twins in the 0.24N alloys compared with the alloys with lower nitrogen concentrations would interrupt the motion of Shockley partial dislocations and enhance the in-grain misorientations. The fracture surfaces of the tensile-tested specimens of the 0N, 0.10N, and 0.24N alloys and their magnified images are shown in Fig. 9a–c and d–f, respectively. Despite their different nitrogen concentrations and different initial microstructures (i.e., volume fractions of the athermal ε martensite phase), all the alloys exhibited intergranular fractures (Fig. 9a–c). In addition, facet planes with striations (or ridges) were observed in all specimens (Fig. 9d–f). These results mean that the fracture mechanisms of all the investigated alloys have the same nature. Since ab initio calculations have shown that a true cleavage fracture does not occur in fcc polycrystalline materials (Morris et al., 2001), the fracture mode in such materials is recognized as quasi-cleavage-type fracture. Takaki et al. (1990) suggested that step-like ridges like those in Fig. 9d–f form through interactions between ε martensite plates, because the primary ε martensite plates (i.e., the ε plates that have been already formed) act as strong obstacles to Shockley partial dislocations moving on the 1 In this study, the KAMs were calculated up to the fifth neighboring point with a maximum misorientation angle of 2˚, in accordance with the procedure described by Herrera et al. (2011).
{111}γ plane along the subsequently formed secondary ε martensite plates (Sato et al., 1982). Therefore, these results suggest that the SIMT generates high amounts of stress that is concentrated at the GBs and/or the γ/ε interfaces, resulting in premature fractures.
4.3. Correlation between the microstructures and the tensile properties 4.3.1.
Strengthening mechanism
In this study, the measured yield stress increased linearly with increasing nitrogen concentration for all alloys except the 0N alloy (Fig. 5a). This was in keeping with previously reported results (Yamanaka et al., 2013), which have suggested that the nanosized precipitates formed when nitrogen is added act as short-range obstacles for gliding Shockley partial dislocations. Nitrogen-induced strengthening in biomedical Co–Cr–Mo alloys has been considered to result from solid solution hardening (Escobedo et al., 1996; Hsu and Lian, 2003). Generally, solid solution hardening of the interstitial elements occupying the octahedral sites in the fcc lattice hardly contributes to the strength, and the interstitial carbon in the Co–29Cr–6Mo alloys considered here actually has a negligible effect on their yield stresses (Lee et al., 2006). Therefore, the remarkable increase in the measured yield stress upon nitrogen addition should be interpreted in terms of precipitation hardening. As mentioned above, the strengthening effect of nitrogen in the Co–29Cr–6Mo alloys (1264 MPa/wt%) is much larger than that observed in austenitic steels ( 400 MPa/wt%; Simmons, 1997). This may also indicate the different strengthening mechanisms in the two alloy systems.
4.3.2.
Ductility
It was seen from Fig. 5b that eliminating the athermal ε martensite from the microstructures is the first step in enhancing the room-temperature ductility of the Co–Cr–Mo alloys. In this
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
process, the nitrogen-induced nanostructural inhomogeneities discussed in this paper hinder the motion of Shockley partial dislocations, which suppresses the athermal martensitic transformation. Adding nitrogen would also influence the kinetics of the SIMT and the resultant work hardening behavior. The work hardening exponent (n) of each alloy was calculated using the following equation (Ludwigson, 1971)2: s ¼ Kεn
ð2Þ
where s and ε are the true stress and true plastic strain during tensile testing, respectively. The results are shown in Table 2. The n values showed a negative correlation with nitrogen concentration, but there was a drastic change in n between the 0.07N and 0.10N alloys. This indicates that nitrogen addition results in a decrease in work hardening that should be interpreted as a retardation of the SIMT. This decrease in the work hardening ability is disadvantageous with respect to tensile elongation, because the plastic instability condition, which is expressed by Eq. (3), is achieved in earlier stages of tensile deformation as the work hardening ability decreases. ds ð3Þ sZ dε On the other hand, the n value is mathematically equal to the true plastic strain at the ultimate tensile strength. Therefore it should be noted that the values of n for all the specimens investigated were larger than the actual tensile elongations and decreased with increasing nitrogen concentration. That is, premature fractures occurred before the onset of plastic instability (i.e., macroscopic necking) in all the Co–29Cr–6Mo alloys investigated. The present study clearly revealed that the formation of the ε martensite must be the origin of such fractures (Fig. 9). Especially for nitrogen concentrations above 0.10 wt%, the elongation-to-failure decreased with increasing nitrogen concentration (Fig. 5b). As shown in Fig. 2, the volume fraction of annealing twins increased upon nitrogen addition. Koizumi et al. (2013) reported the preferential formation of the ε phase produced by the SIMT at ATBs. Similarly, premature fractures related to the action of annealing twins as crack nucleation sites were observed in the Fe–Mn–C austenitic steel (Koyama et al., 2012). Therefore, increasing the amount of annealing twins must result in reduced elongation-to-failure. In summary, we conclude that for Co–29Cr–6Mo alloys, single γ matrices with N concentrations just above 0.10 wt% are preferable for increasing the tensile ductility. This is a novel insight obtained from the results of the current study, because the ductility of the Co–Cr–Mo alloys has previously been considered to increase monotonically with increasing N concentration (Hsu and Lian, 2003; Kilner et al., 1987; Yoda et al., 2012). Co–Cr–Mo alloys with appropriate nitrogen contents may exhibit better self-mating wear behavior, because the nitrogenenhanced strengthening and hardening effects due to the SIMT improve the wear resistance significantly (Chiba et al., 2007). However, since the deformation microstructures in the 0.24N alloys evolved differently than in alloys with lower nitrogen 2 It is well known that Eq. (2) is inadequate for describing the work hardening behavior at low strain levels of fcc alloys with low SFEs, such as austenitic steels. Therefore, the n values were calculated by fitting the flow curves with Eq. (2) at higher strains.
425
concentrations (Fig. 8), the wear behavior of alloys with high nitrogen concentrations will be evaluated in future work.
5.
Conclusions
In the present study, we systematically studied the dependence of the phase constituents, microstructures, tensile properties, and plastic deformation mechanisms of Co– 29Cr–6Mo alloys on the nitrogen concentration. The major findings of the study are as follows:
The volume fraction of the athermal ε (hcp) martensite phase
decreased with increasing nitrogen concentration, and a fully γ (fcc) matrix structure was obtained in the specimens with nitrogen concentrations greater than 0.10 wt%. The addition of nitrogen caused softening of the duplex alloys. This was presumably related to the athermal ε martensite. On the other hand, in specimens with nitrogen contents higher than 0.05 wt%, the yield stress increased monotonically with increasing nitrogen concentration. The addition of nitrogen to biomedical Co–Cr–Mo alloys causes significant strengthening because of nanosized precipitates formed in the γ matrix. Better deformability of the Co–29Cr–6Mo alloys could be obtained by completely suppressing the athermal γ-ε martensitic transformation (at nitrogen contents higher than 0.10 wt%). The further addition of nitrogen slightly decreased the elongation-to-failure, probably because of the enhanced formation of annealing twins, which promoted premature fractures. The addition of nitrogen in excessive amounts, however, reduced the work hardening of the alloys by suppressing the formation of SFs and preventing the SIMT from taking place. This resulted in an intrinsic decrease in the tensile ductility of the alloys.
Acknowledgments The authors would like to thank Isamu Yoshii, Kimio Wako, and Fumiya Sato for sample preparation, and Shun Ito for TEM observations. This research was financially supported by Grant-in-Aid for JSPS Fellows, Global COE Program “Materials Integration (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.
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Available online at www.sciencedirect.com
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Research Paper
Effects of nitrogen addition on microstructure and mechanical behavior of biomedical Co–Cr–Mo alloys Kenta Yamanakaa,n, Manami Morib, Akihiko Chibaa a
Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan NISSAN ARC, LTD., Yokosuka 237-0061, Japan
b
art i cle i nfo
ab st rac t
Article history:
In the present study, the microstructures and tensile deformation behaviors of biomedical
Received 19 July 2013
Co–29Cr–6Mo (wt%) alloys containing different concentrations of nitrogen (0–0.24 wt%)
Received in revised form
were systematically investigated. As the nitrogen concentration increased, the volume
28 September 2013
fraction of athermal ε martensite decreased, because nanoprecipitates hindered the
Accepted 7 October 2013
formation of stacking faults (SFs) by acting as obstacles to Shockley partial dislocation
Available online 23 October 2013
formation, and athermal ε martensite usually forms through the regular overlapping of
Keywords:
SFs. The formation of the athermal ε martensite was completely suppressed when the
Biomedical Co–Cr–Mo alloys
nitrogen concentration exceeded 0.10 wt%, resulting in a simultaneous improvement in
Nitrogen addition
the strength and ductility of the alloys. It was found that the glide of the Shockley partial
Microstructures γ-ε martensitic transformation Mechanical properties
dislocations and the strain-induced γ (fcc)-ε (hcp) martensitic transformation (SIMT) operated as the primary deformation mechanisms. However, adding nitrogen reduced the work hardening by suppressing the formation of the SFs and preventing the SIMT from taking place. This resulted in an intrinsic decrease in the tensile ductility of the alloys. It is also shown that all the alloys exhibited premature fractures owing to the SIMT. The formation of annealing twins in the γ grains is found to be enhanced by nitrogen addition and to promote the SIMT, resulting in a reduction in the elongation-to-failure due to nitrogen addition. These results should aid in the design of alloys that contain nitrogen. & 2013 Elsevier Ltd. All rights reserved.
1.
Introduction
Co–Cr–Mo alloys exhibit excellent biocompatibility, corrosion resistance, and wear resistance, and are therefore widely used for orthopedic implants such as artificial hip and knee joints (Niinomi, 2002; Buford and Goswami, 2004; Chiba et al., 2007). It is known that metal-on-metal bearings made of Co–Cr–Mo alloys allow large-diameter femoral head components to be used, which provide a greater range of motion than that possible with conventional artificial hip joints n
Corresponding author. Tel.: þ81 22 215 2118; fax: þ81 22 215 2116. E-mail address: [email protected] (K. Yamanaka).
1751-6161/$ - see front matter & 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jmbbm.2013.10.006
(Cuckler et al., 2004). Thus, further improvements in the characteristics of these alloys, especially in their wear resistance and mechanical durability, are essential to expand further the biomedical applications of these alloys, and it is of great importance to understand the mechanisms of their functions. In Co–Cr–Mo alloys that comply with the ASTM F75 standard, up to 0.35 wt% carbon can be incorporated to allow for carbide precipitation, which is a major strengthen ing mechanism (Rajan, 1982; Caudillo et al., 2002; Mineta et al., 2010). However, recent studies have revealed that hard
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journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
carbide precipitates are sometimes detrimental to the wear resistance, corrosion resistance, and biocompatibility of the alloys (Chiba et al., 2007; Battini et al., 2011; Liao et al., 2012). Therefore, significant efforts are being made to formu late novel alloys and devise hot deformation processes for these alloys (Salinas-Rodriguez and Rodriguez-Galicia, 1996; Escobedo et al., 1996; Kilner et al., 1987; Hsu and Lian, 2003; Chiba et al., 2007; Matsumoto et al., 2010; Mani, SalinasRodriguez, López, 2011; Yamanaka et al., 2009, 2011, 2012a, 2012b; Mori et al., 2010, 2012; Lee et al., 2008a). The matrix phase also significantly influences the mechanical and tribological properties of Co–Cr–Mo alloys. Generally, these alloys consist of a metastable γ (fcc) matrix and/or platelike ε (hcp) martensite. The ε martensite is formed during quenching, plastic deformation, and isothermal heat treatments and is known to enhance the strength and wear resistance of the alloy (Chiba et al., 2007; Mani et al., 2011). However, the formation of ε martensite also leads to poor deformability (Mori et al., 2010; Mani et al., 2011; Yamanaka et al., 2011, 2012a). Nickel addition is usually employed to prevent such fractures due to the presence of the ε martensite in cobalt-based alloys. In fact, biomedical Co–Cr alloys containing large amounts of Ni (10–37 wt%) (conforming to ASTM F90 and F562) exhibit excellent ductility and have been used in a number of practical applications (Nagai et al., 2012; Marrey et al., 2006). However, nickel is known to cause allergies and cancer in living organisms (Denkhaus and Salnikow, 2002). Nitrogen is known to be an alloying element that can replace nickel. It also stabilizes the γ phase and improves the mechanical properties of Co–Cr–Mo alloys. Early studies have reported that the addition of nitrogen to Co–Cr–Mo alloys results in improvements in the tensile strength, plastic deformability, fatigue strength, and wear properties of the alloys (Escobedo et al., 1996; Hsu and Lian, 2003). Recently, Lee et al. (2008a) and Yoda et al. (2012) investigated the microstructures and tensile properties of as-cast Co–Cr–Mo– N alloys and found that alloys with high nitrogen contents exhibit good combinations of strength and ductility that make them suitable for dental applications. However, the studies performed so far either employed specimens containing relatively large amounts of nickel ( 2.5 wt%) and carbide precipitates (carbon contents 0.45 wt%) (Escobedo et al., 1996; Hsu and Lian, 2003) or varied the chromium concentrations of the samples in order to change their nitrogen concentrations (Lee et al., 2008a; Yoda et al., 2012), which prevents the understanding of the intrinsic effects of nitrogen in Co–Cr–Mo alloys. On the other hand, we have previously reported that Ni- and C-free Co–29Cr–6Mo–N (wt%) alloys exhibit high tensile ductility as well as cold workability while
maintaining their high strength (Lee et al., 2008a; Mori et al., 2010, 2012; Yamanaka et al., 2012b). Nevertheless, a systematic investigation of the effects of nitrogen on the structures and mechanical properties of the Ni- and C-free Co–Cr–Mo biomedical alloys has not been done yet. An important point to note regarding nitrogen-containing Co–Cr–Mo alloys is their decomposition on the nanoscale in the γ phase (Yamanaka et al., 2013). Nanometer-sized nitride precipitates formed in the γ matrix were found to interact with extended dislocations consisting of Shockley partial dislocations and stacking faults (SFs), ultimately inhibiting the γ-ε martensitic transformation (Yamanaka et al., 2013). In this context, the aim of the present study was to investigate the dependence of the microstructures, phase distributions, stabilities of the γ matrices, dislocation structures, and tensile deformation behaviors of Co–29Cr–6Mo–N alloys on the nitrogen concentration.
2.
Materials and methods
2.1.
Specimen preparation
Co–29Cr–6Mo (wt%) alloys with different compositions were prepared in an argon atmosphere using a high-frequency induction furnace. A Cr2N powder was used as the nitrogen source. The chemical compositions of the alloys are shown in Table 1. All the alloys conformed to the ASTM F75 standard and are denoted as 0N, 0.05N, 0.07N, 0.10N, 0.17N, and 0.24N in accordance with the concentration of nitrogen in each alloy. Cast ingots of these alloys of 15 mm in diameter and 200 mm in length were subjected to a homogenizing heat treatment at 1473 K for 1.8 ks and subsequently processed by multipass hot caliber rolling, resulting in a change in their diameter from 15 mm to 9.6 mm. This was followed by the quenching of the hot-rolled alloy bars in water. The equivalent strain, εeq, generated during the rolling process, which was found to be 0.89, was calculated using the following equation: εeq ¼ ln
A0 A
ð1Þ
Here, A0 and A are the cross-sectional areas of the initial and hot-rolled specimens, respectively. Next, the hot-rolled bars were heat treated in air at 1473 K for 0.6 ks and then quenched in water.
Table 1 – Chemical compositions of alloys used in the present study (wt%). Alloy
Co
Cr
Mo
N
Ni
Mn
Si
O
C
0N 0.05N 0.07N 0.10N 0.17N 0.24N
Bal. Bal. Bal. Bal. Bal. Bal.
28.7 27.3 29.0 29.2 27.8 29.9
6.3 6.1 6.0 5.9 6.0 6.1
0.00 0.05 0.07 0.10 0.17 0.24
o0.01 o0.01 o0.01 o0.01 o0.01 o0.01
o0.01 o0.01 o0.01 o0.01 o0.01 o0.01
o0.1 o0.1 o0.1 o0.1 o0.1 o0.1
0.022 0.021 0.021 0.038 0.025 0.028
0.0016 0.0012 0.0012 0.0018 0.0011 0.0028
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
2.2.
Characterization of the microstructures
values were calculated. The fracture surfaces of the alloys were observed using FESEM (Carl Zeiss ULTRA 55).
The microstructures of the alloys were investigated using scanning electron microscopy (SEM), electron-backscatter diffraction (EBSD) analysis, and conventional and high-resolution transmission electron microscopy (TEM and HRTEM, respectively). Cross sections of the specimens, perpendicular to the rolling direction (RD), were observed. The samples subjected to SEM and EBSD analysis were mechanically ground and polished using emery papers and a 0.3 μm alumina suspension. These processes were followed by polishing to a mirror-like finish using a 0.04 μm colloidal silica solution. SEM observations were conducted using a field-emission SEM (FESEM) (Carl Zeiss ULTRA 55) operated at an acceleration voltage of 15 kV. The EBSD measurements were also performed using an FESEM (FEI XL30S-FEG) operated at 15 kV. The EBSD data were collected and analyzed using an orientation image microscope system (TexSEM Laboratories, Inc.). A step size of 0.1 μm in a hexagonal scan grid was employed. TEM and HRTEM were performed using a Topcon EM002B operated at 200 kV. The specimens for TEM were produced by cutting a disk 3 mm in diameter from each alloy and grinding it using a dimple grinder (Gatan Model 656). Next, thin foils were prepared by ion-beam milling these disks (Gatan Model 691, PIPS). To examine the microstructures developed during plastic deformation, TEM observations of the specimens were also performed after they had been subjected to compression tests at a strain rate of 1.0 10 4 s 1 at room temperature, which resulted in a 5% total strain.
2.3.
Tensile tests
Uniaxial tensile tests were performed on the alloy samples at room temperature. Specimens 1.6 1.0 mm2 in the gauge section and 10.5 mm in gauge length were prepared by electrical discharge machining, with the longitudinal axis of the bars parallel to their RD. The samples were strained until failure at a nominal strain rate of 1.6 10 4 s 1. The tensile tests were performed at least three times on each alloy, and the averages and standard deviations (S.D.) of the 0.2% proof stress, ultimate tensile strength, and elongation-to-failure 0.05N
0.07N
3.
Results
3.1.
Microstructures and constituent phases
Fig. 1 shows the EBSD maps of the Co–29Cr–6Mo–N alloys. All the alloys consisted of equiaxed γ grains with a mean grain size of about 100 μm. The phase maps indicate that the 0N, 0.05N, and 0.07N alloys exhibited duplex microstructures comprising a γ matrix and a plate-like ε phase, whereas the 0.10N, 0.17N, and 0.24N alloys exhibited only the γ phase. The ε phase must form through the athermal martensitic transformation occurring during cooling after the heat treatment, because TEM analysis confirmed that both phases have the Shoji–Nishiyama orientation relationship, i.e., (111)γ//(0001)ε and [110]γ//[1120]ε. The inverse pole figure (IPF) and phase maps revealed the formation of multivariant ε martensite structures inside a γ grain, and the number of ε variants decreased with increasing nitrogen content. In the grain boundary (GB) maps, the black, green, and red lines indicate the high-angle boundaries (HABs) with misorientation angles greater than 151, low-angle boundaries (LABs) with misorientation angles of 2–151, and annealing twin boundaries (ATBs) with a Σ3 coincidence site lattice relationship, respectively. The intragranular boundaries in the 0N, 0.05N, and 0.07N alloys were attributable to athermal ε martensite, whereas those in the alloys with over 0.10 wt% nitrogen were due to annealing twins. On the basis of the results of the EBSD analysis, we determined the dependence of the fractions of athermal ε martensite and of ATBs on the nitrogen concentration (Fig. 2). The volume of ε martensite decreased with increasing nitrogen concentration and dropped sharply to zero when the nitrogen concentration was higher than 0.10 wt%. Thus, the critical nitrogen concentration for obtaining a matrix comprising only the γ phase was around 0.10 wt% for the Co–29Cr–6Mo–N system. In contrast, the duplex alloys had 0.10N
0.17N
0.24N
GB
Phase
IPF
0N
100 µm 111 1010 001
101 0001
419
2110
Fig. 1 – EBSD maps of Co–29Cr–6Mo alloys with different nitrogen concentrations.
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journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
Fig. 2 – Volume fractions of athermal ε martensite and ATBs in the Co–29Cr–6Mo alloys as a function of nitrogen concentration.
almost no annealing twins. The volume fraction of ATBs increased rapidly from 0% to 60% for nitrogen concentrations ranging from 0.07 to 0.10 wt%, respectively. In addition, the volume fraction of annealing twins increased gradually from 60% to 75% for nitrogen concentrations ranging from 0.10 to 0.24 wt%, respectively. This was because the formation of multivariant twins within the grains became more pronounced over this range of nitrogen concentrations. Fig. 3a and b shows the SEM–BSE (backscatter electron) images of the 0N and 0.17N alloys with the duplex and single γ matrices, respectively. The microstructure was observed to be subdivided by the athermal ε martensite in Fig. 3a, as well. The s phase, which would appear brighter in these images, was hardly identified in any specimen. Instead, a small number of submicron- and/or micron-sized particles lower in brightness were noted (indicated by the arrow in Fig. 3, for example). The energy dispersive X-ray spectroscopy (EDX) results obtained using the FESEM revealed that these particles were oxides of Cr and Al and might have originated from the raw materials or the alumina crucible used. These oxide particles do not originate from nitrogen addition. They are considered to have a negligible effect on quasi-static tensile deformation behavior discussed below, because their volume fraction is very low.
3.2.
Fig. 3 – SEM–BSE images of (a) 0N and (b) 0.17N alloys subjected to annealing at 1473 K for 0.6 ks, followed by water quenching.
Tensile properties
Fig. 4 shows the nominal stress–nominal strain curves of the Co–29Cr–6Mo–N alloys with nitrogen concentrations ranging up to 0.24 wt%, as obtained from tensile tests. All the specimens were tested to fracture, and macroscopic necking was not noted in any of them. Fig. 5a shows the 0.2% proof stress and ultimate tensile strength of the alloys as a function of their nitrogen concentration (also summarized in Table 2). The 0N alloy exhibited a higher strength than the 0.05N alloy. This may be due to the athermal ε martensite plates, on which the lattice dislocations introduced during water quenching impinge (see Fig. 6a). In contrast, the 0.2% proof
Fig. 4 – Tensile stress–strain curves of Co–29Cr–6Mo alloys with various nitrogen concentrations.
stress was linearly proportional to the nitrogen concentration for concentrations greater 0.05 wt%. It can be seen from Fig. 5a that there was an increase of 1264 MPa in the 0.2% proof stress of the alloys after the incorporation of 1 wt% nitrogen. A similar linear dependence of the strength on the nitrogen concentration has been reported for austenitic steels
421
Elongation (%)
Strength (MPa)
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
Nitrogen content (wt.%)
Nitrogen content (wt.%)
Fig. 5 – Changes in (a) the 0.2% proof stresses and the ultimate tensile strength and (b) uniform elongation as a function of the nitrogen concentration of the Co–29Cr–6Mo–N alloys. The values of the 0.2% proof stress, ultimate tensile strength, and uniform elongation were obtained from the nominal stress–nominal strain curves (Fig. 4). Table 2 – Tensile properties of Co–29Cr–6Mo alloys with different nitrogen concentrations.
0N 0.05N 0.07N 0.10N 0.17N 0.24N
0.2% proof stress (MPa)
Ultimate tensile strength (MPa)
Uniform elongation (%)
Work hardening exponent, n
425758 347724 38176 403736 507715 587734
928747 72874 847722 863741 931711 96279
11.571.9 9.970.6 17.170.3 24.170.5 22.571.2 20.471.1
0.417 0.390 0.384 0.467 0.454 0.405
Fig. 6 – Bright-field TEM images of the Co–29Cr–6Mo alloys with different nitrogen concentrations: (a) 0N and (b) 0.10N alloys.
(Simmons, 1997; Soussan et al., 1991; Lee et al., 2008b). However, the strengthening effect of nitrogen in the present alloys (1264 MPa per wt%) was much more prominent than that in austenitic steels (200–500 MPa per wt%). The ultimate tensile strength decreased and then increased linearly with increasing nitrogen concentration. Fig. 5b shows the dependence of the elongation-to-failure of the Co–29Cr–6Mo–N alloys on the nitrogen content. The 0N alloy exhibited a greater elongation than the 0.05N alloy. On the other hand, it is apparent that eliminating the athermal ε martensite phase (as in alloys with nitrogen contents higher than 0.10 wt%) is an extremely effective way of improving the ductility of the alloys. Since the tensile elongation of the alloys with nitrogen concentrations greater than 0.10 wt%
shows a slight decrease with increasing nitrogen concentration, we can clearly identify the transition area where the elongation-to-failure dramatically changes. The 0.07N alloy falls in this area.
4.
Discussion
4.1. Effect of nitrogen addition on the formation of stacking faults and the athermal γ-ε martensitic transformation Fig. 6a and b shows bright-field TEM images of the 0N and 0.10N alloys, which do and do not contain athermal ε
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Fig. 7 – Bright-field TEM images of the 0.24N alloy annealed at 1473 K for 0.6 ks: (a) before and (b) after compression by 5% at room temperature. HRTEM image taken from the annealed 0.24N alloy specimen is shown in (c).
Table 3 – Equilibrium nitrogen content in the matrix (γ or ε) and Cr2N phases in Co–29Cr–6Mo–(0.1, 0.2)N alloy at various temperatures (wt%) obtained using Thermo-Calc software. Phase
Temperature (K)
Total N content (wt%)
N content in γ or ε phase (wt%)
N content in Cr2N phase (wt%)
Mole fraction of Cr2N phase (%)
γ phase
1473 1473 1273 1273
0.1 0.2 0.1 0.2
8.51 10 3 8.45 10 3 1.10 10 3 1.10 10 3
10.4 10.5 11.0 10.9
1.20 2.49 1.26 2.53
ε phase
1073 1073 873 873
0.1 0.2 0.1 0.2
8.25 10 3 8.25 10 3 0.67 10 3 0.67 10 3
11.3 11.3 11.7 11.7
1.17 2.42 1.24 2.48
martensite, respectively. Dislocations in both alloys were found to have frequently dissociated into SFs, forming highly planar dislocation arrays. Compared with the nitrogenbearing austenitic steels (Ojima et al., 2009), both alloys have very high widths of their SFs. These dislocation structures indicate that the stacking fault energies (SFEs) of the alloys are very low, as predicted by thermodynamic calculations; the SFE values eventually become negative at temperatures lower than 1100 K (Yamanaka et al., 2009, 2012a). Although the SFEs are usually positive for conventional metals and alloys, a negative SFE simply means that the ε phase is more stable than the γ phase. It is well known that SFs can be regarded as local hcp structures embedded in the fcc matrix. Thus, the formation of thermodynamically stable SFs and ε martensite is preferred in the Co–29Cr–6Mo alloys. A similar dislocation structure was also observed in the alloy containing 0.20 wt% N in our previous study (Yamanaka et al., 2013). The addition of nitrogen in greater amounts resulted in significantly different states of the dislocations than those present in the alloys with lower nitrogen concentrations. Fig. 7a shows bright-field TEM images of the 0.24N alloys. Although the black contrast in this figure corresponds to the pairs of Shockley partial dislocations bounding SFs, the width of SFs is not as large as in Fig. 6. In contrast, Fig. 7b shows a bright-field TEM image of the dislocation structures in the 0.24N alloy deformed by 5% during compression at room temperature. It should be noted that the dissociation of the dislocations increased after the deformation and resulted in
the formation of a planar dislocation structure consisting of widely extended SFs. This structure is similar to those of the alloys with lower nitrogen concentrations (Fig. 6). The transition of dislocation structures observed here should be understood in light of nanoprecipitate formation in the γ matrix originating from the nitrogen doping. An HRTEM image of the 0.24N alloy is shown in Fig. 7c. The corresponding selected area diffraction (SAD) pattern for an incident beam parallel to the [110]γ direction is superimposed in the figure. By analyzing the SAD patterns obtained under different incident beam conditions, we identified nanosized plate-like Cr2N precipitates that formed on the {111} planes of the γ phase. Furthermore, streaks originating from the precipitates were identified along the o1114γ directions in the SAD pattern. The nitrogen contents in the γ, ε, and Cr2N phases as a function of temperature and total nitrogen content of the alloys were calculated using the Thermo-Calc software (Table 3). The thermodynamic data sets needed for this calculation were acquired from TCS Steels/Fe-alloys Database Version 6. Note that the ε phase is stable below 1200 K in the Co–29Cr–6Mo–N alloys (Yamanaka et al., 2013). The thermodynamic calculations indicated that the solubilities of nitrogen in the γ and ε phases are both quite low, and therefore the Cr2N phase forms in both the γ and ε phases. In addition, the N concentrations in the Cr2N precipitates are determined by the temperature and are almost independent of the total N content. In other words, increasing the bulk nitrogen content results in an increase in the fraction of Cr2N particles. In addition, the amount of Cr2N phase is
journal of the mechanical behavior of biomedical materials 29 (2014) 417 –426
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Fig. 8 – EBSD maps of the microstructures near the fracture surfaces of the (a–d) 0.10N, and (e–h) 0.24N alloys deformed by tensile stress to fracture: (a and e) IPF, (b and f) phase, (c and g) GB, and (d and h) KAM maps.
expected to increase with increasing total N content. The previous study (Yamanaka et al., 2013) revealed that the Cr2N nanoprecipitates act as obstacles for the Shockley partial dislocations and ultimately retard the γ-ε martensitic transformation. Since the amount of the Cr2N phase increases with increasing N concentration (Table 3), the volume fraction of athermal ε martensite consequently decreased with increasing N concentration (Fig. 2). These nanoprecipitates can be overcome under the combined action of the applied stress and thermal energy. Therefore, wide SFs can form even in the 0.24N alloys when a sufficiently high stress is applied (Fig. 7b). Namely, these results indicate that adding nitrogen reduces the planarity in dislocation structures of the Co–Cr–Mo alloys, but the width of SFs becomes larger during tensile deformation and, consequently, the highly planar dislocation structure forms even in the high-nitrogen alloys. Note that perfect dislocations as well as dislocation cells/walls are hardly observed in our alloys investigated in this research, because the SFE of the alloys is negative (Yamanaka et al., 2009, 2012a).
4.2.
Plastic deformation and fracture mechanisms
Fig. 8 shows EBSD maps of cross sections of the tensilefractured specimens of the (a–d) 0.10N and (e–h) 0.24N
alloys. As mentioned above, both alloys contain a negligibly small amount of the athermal ε martensite phase (Fig. 1) and exhibited similar elongations-to-failure (Table 2). In the IPF maps (Fig. 8a and e) and phase maps (Fig. 8b and f) of both alloy specimens, we can identify the in-grain orientation gradient and the ε phase, which is formed via the strain-induced γ-ε martensitic transformation (SIMT) during room-temperature deformation. Thus, for the nitrogen concentrations investigated in this study up to 0.24 wt%, the deformation mode was slipping of the Shockley partial dislocations and subsequent formation of ε martensite, although Yoda et al. (2012) have stated that adding nitrogen changes the deformation mechanism from the SIMT to another form such as twinning or dislocation slipping. Despite these similar IPF and phase maps, the GB maps revealed that quite different microstructures developed during tensile deformation of the two alloys: there were intragranular HABs related to the SIMT in the 0.10N alloys (Fig. 8c), whereas a significant number of deformation-induced LABs developed in the 0.24N alloys (Fig. 8g). Accordingly, the distribution of the local strain was different in the two alloys. The kernel average misorientation (KAM) maps of the alloys are shown in Fig. 8d and h. The KAM value represents an average misorientation angle between all adjacent
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Fig. 9 – SEM images showing typical fracture surfaces of the Co–29Cr–6Mo alloys with different nitrogen concentrations: (a and b) 0N, (c and d) 0.10N, and (e and f) 0.24N alloys.
measurement points within a grain1 and is correlated with the density of geometrically necessary dislocations (Calcagnotto et al., 2010). In the 0.10N alloys (Fig. 8d), the KAM values were preferentially higher in the vicinities of the γ GBs and in the ε martensite plates. These results can be interpreted in terms of GB compatibility and impingement of the dislocations or the plate-like ε phase formed by the SIMT at these boundaries. In contrast, the 0.24N alloys exhibited a more uniform distribution of KAMs inside the γ grains (Fig. 8h). The higher density of Cr2N nanoprecipitates and higher fraction of annealing twins in the 0.24N alloys compared with the alloys with lower nitrogen concentrations would interrupt the motion of Shockley partial dislocations and enhance the in-grain misorientations. The fracture surfaces of the tensile-tested specimens of the 0N, 0.10N, and 0.24N alloys and their magnified images are shown in Fig. 9a–c and d–f, respectively. Despite their different nitrogen concentrations and different initial microstructures (i.e., volume fractions of the athermal ε martensite phase), all the alloys exhibited intergranular fractures (Fig. 9a–c). In addition, facet planes with striations (or ridges) were observed in all specimens (Fig. 9d–f). These results mean that the fracture mechanisms of all the investigated alloys have the same nature. Since ab initio calculations have shown that a true cleavage fracture does not occur in fcc polycrystalline materials (Morris et al., 2001), the fracture mode in such materials is recognized as quasi-cleavage-type fracture. Takaki et al. (1990) suggested that step-like ridges like those in Fig. 9d–f form through interactions between ε martensite plates, because the primary ε martensite plates (i.e., the ε plates that have been already formed) act as strong obstacles to Shockley partial dislocations moving on the 1 In this study, the KAMs were calculated up to the fifth neighboring point with a maximum misorientation angle of 2˚, in accordance with the procedure described by Herrera et al. (2011).
{111}γ plane along the subsequently formed secondary ε martensite plates (Sato et al., 1982). Therefore, these results suggest that the SIMT generates high amounts of stress that is concentrated at the GBs and/or the γ/ε interfaces, resulting in premature fractures.
4.3. Correlation between the microstructures and the tensile properties 4.3.1.
Strengthening mechanism
In this study, the measured yield stress increased linearly with increasing nitrogen concentration for all alloys except the 0N alloy (Fig. 5a). This was in keeping with previously reported results (Yamanaka et al., 2013), which have suggested that the nanosized precipitates formed when nitrogen is added act as short-range obstacles for gliding Shockley partial dislocations. Nitrogen-induced strengthening in biomedical Co–Cr–Mo alloys has been considered to result from solid solution hardening (Escobedo et al., 1996; Hsu and Lian, 2003). Generally, solid solution hardening of the interstitial elements occupying the octahedral sites in the fcc lattice hardly contributes to the strength, and the interstitial carbon in the Co–29Cr–6Mo alloys considered here actually has a negligible effect on their yield stresses (Lee et al., 2006). Therefore, the remarkable increase in the measured yield stress upon nitrogen addition should be interpreted in terms of precipitation hardening. As mentioned above, the strengthening effect of nitrogen in the Co–29Cr–6Mo alloys (1264 MPa/wt%) is much larger than that observed in austenitic steels ( 400 MPa/wt%; Simmons, 1997). This may also indicate the different strengthening mechanisms in the two alloy systems.
4.3.2.
Ductility
It was seen from Fig. 5b that eliminating the athermal ε martensite from the microstructures is the first step in enhancing the room-temperature ductility of the Co–Cr–Mo alloys. In this
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process, the nitrogen-induced nanostructural inhomogeneities discussed in this paper hinder the motion of Shockley partial dislocations, which suppresses the athermal martensitic transformation. Adding nitrogen would also influence the kinetics of the SIMT and the resultant work hardening behavior. The work hardening exponent (n) of each alloy was calculated using the following equation (Ludwigson, 1971)2: s ¼ Kεn
ð2Þ
where s and ε are the true stress and true plastic strain during tensile testing, respectively. The results are shown in Table 2. The n values showed a negative correlation with nitrogen concentration, but there was a drastic change in n between the 0.07N and 0.10N alloys. This indicates that nitrogen addition results in a decrease in work hardening that should be interpreted as a retardation of the SIMT. This decrease in the work hardening ability is disadvantageous with respect to tensile elongation, because the plastic instability condition, which is expressed by Eq. (3), is achieved in earlier stages of tensile deformation as the work hardening ability decreases. ds ð3Þ sZ dε On the other hand, the n value is mathematically equal to the true plastic strain at the ultimate tensile strength. Therefore it should be noted that the values of n for all the specimens investigated were larger than the actual tensile elongations and decreased with increasing nitrogen concentration. That is, premature fractures occurred before the onset of plastic instability (i.e., macroscopic necking) in all the Co–29Cr–6Mo alloys investigated. The present study clearly revealed that the formation of the ε martensite must be the origin of such fractures (Fig. 9). Especially for nitrogen concentrations above 0.10 wt%, the elongation-to-failure decreased with increasing nitrogen concentration (Fig. 5b). As shown in Fig. 2, the volume fraction of annealing twins increased upon nitrogen addition. Koizumi et al. (2013) reported the preferential formation of the ε phase produced by the SIMT at ATBs. Similarly, premature fractures related to the action of annealing twins as crack nucleation sites were observed in the Fe–Mn–C austenitic steel (Koyama et al., 2012). Therefore, increasing the amount of annealing twins must result in reduced elongation-to-failure. In summary, we conclude that for Co–29Cr–6Mo alloys, single γ matrices with N concentrations just above 0.10 wt% are preferable for increasing the tensile ductility. This is a novel insight obtained from the results of the current study, because the ductility of the Co–Cr–Mo alloys has previously been considered to increase monotonically with increasing N concentration (Hsu and Lian, 2003; Kilner et al., 1987; Yoda et al., 2012). Co–Cr–Mo alloys with appropriate nitrogen contents may exhibit better self-mating wear behavior, because the nitrogenenhanced strengthening and hardening effects due to the SIMT improve the wear resistance significantly (Chiba et al., 2007). However, since the deformation microstructures in the 0.24N alloys evolved differently than in alloys with lower nitrogen 2 It is well known that Eq. (2) is inadequate for describing the work hardening behavior at low strain levels of fcc alloys with low SFEs, such as austenitic steels. Therefore, the n values were calculated by fitting the flow curves with Eq. (2) at higher strains.
425
concentrations (Fig. 8), the wear behavior of alloys with high nitrogen concentrations will be evaluated in future work.
5.
Conclusions
In the present study, we systematically studied the dependence of the phase constituents, microstructures, tensile properties, and plastic deformation mechanisms of Co– 29Cr–6Mo alloys on the nitrogen concentration. The major findings of the study are as follows:
The volume fraction of the athermal ε (hcp) martensite phase
decreased with increasing nitrogen concentration, and a fully γ (fcc) matrix structure was obtained in the specimens with nitrogen concentrations greater than 0.10 wt%. The addition of nitrogen caused softening of the duplex alloys. This was presumably related to the athermal ε martensite. On the other hand, in specimens with nitrogen contents higher than 0.05 wt%, the yield stress increased monotonically with increasing nitrogen concentration. The addition of nitrogen to biomedical Co–Cr–Mo alloys causes significant strengthening because of nanosized precipitates formed in the γ matrix. Better deformability of the Co–29Cr–6Mo alloys could be obtained by completely suppressing the athermal γ-ε martensitic transformation (at nitrogen contents higher than 0.10 wt%). The further addition of nitrogen slightly decreased the elongation-to-failure, probably because of the enhanced formation of annealing twins, which promoted premature fractures. The addition of nitrogen in excessive amounts, however, reduced the work hardening of the alloys by suppressing the formation of SFs and preventing the SIMT from taking place. This resulted in an intrinsic decrease in the tensile ductility of the alloys.
Acknowledgments The authors would like to thank Isamu Yoshii, Kimio Wako, and Fumiya Sato for sample preparation, and Shun Ito for TEM observations. This research was financially supported by Grant-in-Aid for JSPS Fellows, Global COE Program “Materials Integration (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.
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