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Effect of thermo-mechanical treatment on superelastic behavior of Ti–19Nb–14Zr (at.%) shape memory alloy
Intermetallics 32 (2013) 44e50
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Effect of thermo-mechanical treatment on superelastic behavior of Tie19Nbe14Zr (at.%) shape memory alloy L.W. Ma, H.S. Cheng, C.Y. Chung* Department of Physics and Materials Science, City University of Hong Kong, 83 Tat Chee Avenue, Kowloon Tong, Hong Kong
a r t i c l e i n f o
a b s t r a c t
Article history: Received 16 February 2012 Received in revised form 29 May 2012 Accepted 19 July 2012 Available online 9 October 2012
The effect of thermo-mechanical treatment on shape memory effect and superelastic behavior of Ti e19Nbe14Zr (at.%) shape memory alloy is investigated. The annealed specimens are aged at 300 C for different times from 10 min to 20 h followed by ice water quenching. Both reverse and forward martensitic transformation peaks are identified in the annealed and aged specimens. Fine u particles with a dimension of w20 nm in b grains are clearly observed in the TEM micrograph. The strains of the annealed specimen are completely recovered by heating, starting from the 2nd cycle to the 5th cycle due to the reverse martensitic transformation. Non-martensitic superelasticity strain of about 6% is obtained in the aged-10 min specimen. It is noted that the u phase hinders the formation of stress-induced martensite from the b phase, suggesting that the non-martensitic superelasticity should be attributed to the presence of u þ b phases. Also, the non-martensitic superelasticity can be improved through increasing the critical stress for dislocation slip deformation by controlling the size and amount of u phase and/or introducing plastic strain. It is found that this non-martensitic superelasticity can be obtained from the present alloy with combined u þ b phases by introducing additional precipitation hardening and/or work hardening. The elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately of 14 GPa, which is close to that of cortical bone of human beings (7 e35 GPa). Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: B. Martensitic transformations B. Shape-memory effect C. Thermomechanical treatment F. Mechanical testing
1. Introduction Titanium alloys have been widely used for biomedical implants because of their superior corrosion resistance, mechanical properties and tissue acceptance characteristics. However, many studies have reported that unsatisfactory loads transfer from the implant devices and the relatively high elastic modulus of implant materials may lead to bone resorption [1]. The elastic moduli mismatch between the metallic implant and adjacent bone is often referred to the cause of this ‘stress shielding effect’ [2]. Developing alloys with low modulus and superelasticity may provide solutions to ease the stress shielding problem [3,4]. TieNi alloys have been applied to various medical products due to their excellent superelasticity under room and body temperature [5e7]. However, it has been reported that pure Ni is toxic and may cause Nihypersensitivity [8]. Hence, new alloys which are Ni-free have been developed for biomedical applications to substitute TieNi alloys. TieNb based shape memory alloy is one of the candidates which have recently attracted considerable attention for their potential to * Corresponding author. Tel.: þ852 3442 7835; fax: þ852 3442 0538. E-mail address: [email protected] (C.Y. Chung). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.07.024
be used in biomedical devices. It is because they contain no toxic elements and possess superior mechanical properties [9e11]. Many researches have investigated the superelasticity of TieNb based alloys [12,13]. However, the reported superelasticity characteristic is not as good as that of TiNi. Moreover, the incomplete superelastic behavior was unsatisfactory. It is well known that the shape memory effect and superelasticity is based on the reversible b / a00 transformation at loading temperature higher than Af, superelasticity is limited to alloys with Af lower than the temperature of the loading temperature. The transformation peak corresponding to martensitic transformation in TieNb was difficult to be identified using DSC due to small enthalpy of transformation and the large difference between Ms and Mf [12]. Hence, the study and application of Ni-free SMA in human body is not straightforward. Omega (u) phase is a metastable phase in Ti-based alloys with hexagonal structure. It can be easily formed either by quenching from high temperature b phase (athermal-uath) or by aging at intermediate temperatures (isothermal-uiso) [14,15]. It had been reported that the precipitation of u phase had caused embrittlement [16,17]. On the other hand, some papers had suggested that fine u precipitates are effective in improving the superelastic
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property of the titanium alloy [18,19]. It is quite obvious that the presence of u phase in shape memory alloy would affect the shape memory effect and/or superelasticity of titanium alloys; however, it is not well understood. The objective of the present work is to specifically study the effects of thermo-mechanical treatment on superelastic behavior of Tie19Nbe14Zr (at.%) shape memory alloy. The transformation temperatures of the alloy subjected to annealing and different aging time were determined using the Differential Scanning Calorimetry (DSC) thermal analysis. X-ray diffraction (XRD) was used to examine the crystallographic structures of the annealed and aged specimens. Besides, transmission electron microscopy (TEM) was used to reveal the presence of u phase in the aged specimen. Finally, the shape memory effect and superelastic behavior of specimens subjected to both annealing and aging were also investigated by cyclic tensile test. 2. Material and methods The alloy ingot with Tie19Nbe14Zr (at.%) was fabricated by a vacuum consumable arc melting furnace. The ingot was coldrolled to a plate with a reduction in thickness of 90%. Specimens for DSC analysis, XRD measurement, TEM observation and tensile tests were cut from the plate by an electro-discharge machine. The specimens were annealed at 600 C for 30 min followed by ice water quenching. Some of the annealed specimens were then aged at 300 C from 10 min to 20 h followed by ice water quenching. Hereinafter, annealed specimens followed by aging at 300 C with different time-x are abbreviated to aged-x specimens, where x ranged from 10 min to 20 h. An aqueous solution of HF and HNO3 was used to remove the oxidized layer of both the annealed and aged specimens. Phase transformation temperatures were determined between 100 C and 600 C using a Perkin Elmer (PE) Differential Scanning Calorimetry (DSC7) thermal analysis. The specimen weight for DSC thermal analysis was w50 mg, while the heating and cooling rate was 20 K/min. The transformation start and finish temperatures were determined by the intersection of a base line and the tangent to a peak in the DSC curves. The absorbed/released heats during the transformations were calculated from the areas under curves of the heat flows between the start and finish temperatures of transformations. The constituent phases of the specimens were identified using an X’pert MPD Pro XA) ray diffraction (XRD) instrument with Cu-Ka target (l ¼ 1.5406 in 2q ranges between 30 and 80 . TEM specimens were prepared by electrochemical polishing using an electrolyte of 6% perchloric acid þ 35% butanol þ 59% methanol at 10 C. TEM observations were performed using a Philips CM20 microscope. Tensile cycling test and training were carried out at a strain rate of 1.9 104 s1 at room temperature. The gage length of the specimens for tensile tests was 14 mm. Finally, the stress-induced martensite variants of aged specimen after cyclic tensile test were revealed using a JEOL JSM-820 Scanning Electron Microscope (SEM). 3. Results and discussion The transformation profiles of the specimens subjected to different heat treatments were measured by the DSC analysis. All specimens were initially heated to w500 C and then cooled down to w100 C. Fig. 1 shows the DSC curves of (a) the annealed specimen, (b) the aged-1 h specimen and (c) the aged-20 h specimen. All heating and cooling curves were represented using solid lines and dash lines respectively. Two distinct peaks were detected in all specimens. These two peaks are proved a pair of transformation using the “partial DSC cycle technique”, which was usually performed to identify the nature of multi-stage transformation [20]. It
Fig. 1. DSC curves of (a) the annealed specimen, (b) the aged-1 h specimen and (c) the aged-20 h specimens.
is believed that a pair of peaks in each specimen represents the reverse and forward martensitic transformations. Further clarification will be explained by shape memory recovery using stresse strain curves. A slight shift in peak of reverse martensitic transformation was observed in the aged specimens. It is found that both As and Af increase with increase in aging time. It is also noted that the absorbed energy (Eabs) during reverse martensitic transformation increases with longer aging time. The increase in Af and Eabs is possibly due to increasing the stability of martensite by martensitic aging effect [21,22] and/or precipitation of u phase. In contrast, the Ms decreased very slightly with longer aging time. These changes in Ms may be attributed to the precipitation of u phase and/or parent phase aging [23]. Isothermal omega phase (uiso) was reported to have precipitated in b-Ti based alloys when subjected to subsequent aging effect [24]. Subsequently, the formation of uiso may lead to an increase in the b-stabilizer in the matrix, causing the Ms to decrease. The transformation temperatures and energies of the annealed, aged-1 h and aged-20 h specimens obtained from the DSC curves are listed in Table 1. It should be pointed out that the martensitic finish temperatures of all samples are above room temperature. This means that the present alloys should be in martensite state at room temperature. Fig. 2 shows the XRD patterns of the (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen, and (d) aged-20 h specimen. In Fig. 2(a), b phase was surprisingly found to be the predominant phase with the presence of a small amount of u and a00 phases (annealed specimen). Since the Ms and Mf of annealed specimen was measured to be above room temperature (Table 1), expecting that the annealed specimen followed by quenching comprises primarily of a00 martensite rather than b phase. However, only a very weak peak corresponding to a00 martensite was detected. It is
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Table 1 Transformation temperatures and energies of the annealed, aged-1 h and aged-20 h specimens.
Annealed specimen Aged-1 h specimen Aged-20 h specimen
Mf
Mp
Ms
Erel
Af
Ap
As
Eabs
138 156 146
212 209 192
250 246 242
2.2 1.8 1.6
419 428 438
384 400 409
310 331 336
14.9 22.6 30.5
Where Mf ¼ martensite finish temperature ( C), Mp ¼ martensite peak temperature ( C), Ms ¼ martensite start temperature ( C), Erel ¼ released energy (J/g), Af ¼ austenite finish temperature ( C), Ap ¼ austenite peak temperature ( C), As ¼ martensite start temperature ( C), Eabs ¼ absorbed energy (J/g).
suggested that the formation of small amount of uath phase could suppress a large amount of transformations from b to a00 , although forward martensitic transformation was observed above room temperature in the DSC curve (Fig. 1(a)). The present observations are similar to those reported by Zain et al. [25] and Lin et al. [26], who investigated the superelastic behavior of TieNbeMo alloys, and the structure and properties of Ti-7.5Mo-xFe alloys respectively. Both groups also found that the formation of small amount of uath phase after quenching suppressed the martensitic transformation. Also, a second stage martensitic transformation [27] may be another possible explanation for the existence of b phase below Mf temperature as a major phase component. For further confirmation, the annealed specimen was initially immersed by liquid nitrogen and then characterized by the XRD analysis. However, the XRD results showed no significant difference with that in Fig. 2(a), the remaining b phase maintained ‘frozen’ or untransformed even at a temperature of about w200 C (liquid nitrogen). Nevertheless, similar ‘frozen’ b structure has been reported in slightly doped Ti50xNi50þx [28] and Ti50(Pd50xCrx) [29] alloys. They both claimed that introducing excess dopants to a shape memory alloy results in a random distribution of a small amount of quasistatic local strain domains or nanodomains which decreases the thermodynamic driving force for forming martensite. The growth process of martensite domains can also be retarded by the pre-existing quasistatic nanodomains. Therefore, b phase was found to be stable at room temperature. In Fig. 2(b), similar XRD spectrum was obtained in the aged10 min specimen as compared with the annealed specimen. However, the volume fraction of u phase increased with decreased in peak intensity of b phase in the aged-1 h specimen (Fig. 2(c)). It is possibly due to the transformation from b to uiso during aging, since the u phase was reported to be easily formed by either quenching (athermal) or aging (isothermal). Interestingly, the volume fraction of both a00 and u phases were simultaneously increased when the
Fig. 2. XRD patterns of the (a) annealed specimen, (b) aged-10 min specimen, (c) aged1 h specimen, and (d) aged-20 h specimen.
specimen aged for 20 h (Fig. 2(d)). It is clear that longer aging time can increase the volume fraction of uiso phase [18,19,30]. Therefore, the specimen aged for 20 h consistently has the largest content of uiso phase. Also, the energy of forward martensitic transformation was decreased from 2.2 to 1.6 J/g, suggesting the uiso phase suppressed a part of martensitic transformation by longer aging time. However, the volume fraction of a00 phase also increased with longer aging time should be associated with the transformation from uath back to a00 martensite. Hence, it is suggested that there are two transformations, b / uiso and uath / a00 , occurred in the aged specimen for 20 h. This is in agreement with the phase transformation of a00 martensite structure in aged Tie8wt%Mo alloy reported by Mantani and co-workers [21]. They reported that a00 twins were formed when the u phase products were changed into the a00 martensite by aging. Fig. 3 shows (a) the bright-field TEM image of the aged-1 h specimen, (b) the corresponding selected area diffraction pattern (SADP) of Fig. 3(a) and (c) the key diagram of diffraction pattern. Fine u particles with a dimension of w20 nm in b grains were clearly observed. The selected diffraction pattern was obtained from the [110]b zone axis. The diffraction pattern of the b þ u matrix is consistent with the work of Kim et al. [18]. Strong u reflections in the diffraction pattern indicated that there is a high volume fraction of u phase in the selected area. Also, extra reflections resulting from double diffraction [31] were observed in the pattern, it has been reported in previous studies that the presence of u phase in b grains was always accompanied with broadening of diffraction spots and diffuse streaks [32e34] resulting from the incomplete (111) plane collapse from b phase. Tensile tests were performed at room temperature to study the effect of thermo-mechanical process to shape memory properties of the specimens. Fig. 4 shows the stressestrain curves of (a) asrolled specimen, (b) annealed specimen (c) aged-10 min specimen and (d) aged-1 h specimen. The as-rolled specimen exhibit a single stage yielding with only small plastic deformation and high critical stress compared with other heat-treated specimens. It is due to the work hardening effect induced by cold rolling. Both of the annealed (Fig. 4(b)) and aged10 min (Fig. 4(c)) specimens exhibit a two-stage yielding [12]. It should be noted again that b phase was observed as a predominant phase in all specimens according to the XRD patterns shown in Fig. 2. Typically, stress induced martensitic transformation occurs at a temperature above Af in b phase. However, the b phase is ‘frozen’ even at a temperature below Mf (i.e. room temperature), it is believed stress induced martensitic transformation also occurs in this frozen b structure at room temperature [28,29]. Hence, the apparent yielding stress (first stage yielding) mainly corresponds to the critical stress for inducing the martensitic transformation upon loading (sSIM). Since the As of the alloy is much higher than the loading temperature (i.e. room temperature) as detected by DSC, the stress-induced martensites were maintained. The slope after the apparent yielding stress is related to the transformation of stress-induced martensite (MSI) to martensite variants (Mv) (second stage yielding). In Fig. 4(c), the slope for the transformations of MSI / Mv and critical stress for slip increased after aging as compared with the annealed specimen (Fig. 4(b)). The ductility of the aged-10 min specimen was decreased by this hardening effect and it is apparently due to the uiso precipitation as confirmed by the XRD profiles or the structural change of a00 martensite [21]. However, the aged-1 h specimen did not exhibit two-stage yielding, though there is minor difference in the XRD profiles between the annealed and the aged-1 h specimens. Furthermore, the critical stress of the aged-1 h specimen is the highest one among all specimens. It is suggested that the uiso precipitates are effective in increasing the critical stress for forming
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Fig. 3. (a) The bright-field TEM image of the aged-1 h specimen, (b) the corresponding selected area diffraction pattern (SADP) of Fig. 3(a) and (c) the key diagram of diffraction pattern.
stress-induced martensite and the stress for the transformation of M / Mv. Therefore, no significant amount of stress-induced martensite and martensite variants were formed before plastic deformation. It is also found that the elastic modulus of the specimen is sensitive to the aging time. The elastic modulus of the annealed specimen increases from 14 GPa to 23 GPa after aging for 1 h. This indicates that the elastic modulus increases with
Fig. 4. Shows the stressestrain curves of (a) as-rolled specimen, (b) annealed specimen (c) aged-10 min specimen and (d) aged-1 h specimen.
increasing the volume fraction of u phase as the Young’s modulus of u phase is much higher than that of b and a00 phase [35]. To further investigate the shape memory and superelastic properties of the alloy, cyclic loading and unloading tensile test was performed at the room temperature. Fig. 5 shows the stressestrain curves obtained by cyclic loadingeunloading tensile test of (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen and (d) the aged-10 min specimen after five cycles of loadingeunloading tests in Fig. 5(c). The annealed specimen was loaded up to 2% in strain and then unloaded during the first cycle, Fig. 5(a). The measurement was repeated by increasing the maximum strain by 1% upon loading for the same annealed specimen followed by heating to a temperature higher than Af, i.e. 450 C, after each cycle. The shape memory recovery and superelasticity were confirmed in the annealed specimen, a line with an arrow indicates the shape recovery by heating. The shape memory recovery of the annealed specimen (Fig. 5(a)) occurred only when the specimen reached a temperature of about 450 C which is consistent with the peak on the heating curve of DSC result (Fig. 1(a)), hence it is believed that the peak on the heating curve is due to reverse martensitic transformation while the peak obtained during cooling corresponds to the forward martensitic transformation. The strains are completely recovered by heating, starting from the 2nd cycle to the 5th cycle due to the reverse martensitic transformation. The strains obtained in the curves can be defined into three types: (1) the elastic strain (3 el) recovered
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Fig. 5. The stressestrain curves obtained by cyclic loadingeunloading tensile test for (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen and (d) the aged10 min specimen which already subjecting to five cycles of loadingeunloading tests in Fig. 5(b).
elastically upon unloading, (2) the strain recovered superelastically upon unloading (3 se) and (3) the strain recovered by heating (3 sme). After five cycles, the annealed specimen exhibits complete shape memory effect and superelasticity up to the maximum strain of w8%. Hence the sum of 3 se, 3 el, 3 sme as shown in Fig. 5(a) can be indicated as the total recoverable strain (3 r). It should be noted that the superelasticity of SMA is usually related to stress-induced reversible martensitic transformation. In the present study, the superelasticity (non-martensitic) obtained was apparently not associated with the stress-induced reversible martensitic transformation since the reverse martensitic transformation temperatures of the present alloy were confirmed to be much higher than room temperature by DSC analysis (Fig.1). It was reported that “nonmartensitic” alloy Ni48.5Ti51.5 can exhibit both SME and SE without showing martensitic transformation [36]. Furthermore, nonmartensitic superelasticity can also be resulted from other kinds of mechanisms, such as dislocation based mechanism of incipient kink band (IKB) [37] and dislocation free mechanism of giant fault [38] or dislocation-free sliding. In Fig. 5(b), almost complete nonmartensitic superelasticity was observed in aged-10 min specimen during the first and second cycles. However, only single stage yielding was obtained in aged-1 h specimen (Fig. 5(c)). It is suggested that the critical stress for inducing martensite is much higher than the critical stress for slip deformation. A severe loss of ductility was obtained in the aged-1 hr. specimen compared with the aged-10 min specimen. This implies that the precipitations of more thermal u phase formed after longer aging time is effective in hardening. Hence, it is considered that Tie 19Nbe14Zr is not a promising high temperature shape memory alloy as it leads to brittleness when subjecting at high temperature. In contrast, an extra cyclic loading and unloading test (Fig. 5(d)) was specifically performed on the same aged-10 min specimen which already subjecting to five cycles of loadingeunloading test in Fig. 5(c). Complete non-martensitic superelasticity with approximately of 6% was obtained.
In this investigation, the actual mechanism for the “nonmartensitic superelasticity” obtained in the present alloy should be associated with the combination of u þ b phases and/or the atomic rearrangement within the same sub-lattice of the imperfectly ordered alloy during martensite aging [22]. It is suggested that the u phase hinders the formation of stress-induced martensite from the b phase. Hence, the non-martensitic superelasticity can be improved in this combination of u þ b phases with extra appropriate precipitation hardening (aging) and/or strain hardening. The critical stress for slip deformation (sCSS), the critical stress for inducing martensite (sSIM) and the elastic modulus (E) for aged10 min specimen estimated in Fig. 5(b) was plotted against the maximum strain (%) in Fig. 6.
Fig. 6. The critical stress for slip deformation (sCSS), the critical stress for inducing martensite (sSIM) and the elastic modulus (E) for aged-10 min specimen estimated in Fig. 5(b) was plotted against the maximum strain (%).
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Fig. 7. SEM micrographs of the aged-10 min specimen (a) before the cyclic tensile test, and (b) after the cyclic test with maximum strain of 8%.
The critical stress for slip deformation (sCSS) increased continuously, should be associated with the work hardening effect during mechanical cycling. However, the critical stress for inducing martensite (sSIM) decreased with increasing the maximum loading strain. The elastic modulus is approximately of 14 GPa and decreases slightly with increasing maximum strains. The increase in volume fraction of residual a00 phase can be attributed to stress induced martensitic transformation upon loading and unloading. SEM examination was used to study the structural change of the alloy after the cyclic tensile test. Fig. 7 shows the SEM micrographs of the aged-10 min specimen (a) before the cyclic tensile test, and (b) after the cyclic test with maximum strain of 8%. Stress-induced martensite variants with lamellar structures were observed within the equi-axed b grains. It is clear that there is a transformation of almost entire b phase to a b þ a00 system due to the formation of stress induced martensite. Also, the elastic modulus of the aged-10 min specimen decreased slightly upon cyclic loading due to the lower modulus nature of the a00 phase [39,40]. The elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately of 14 GPa, which is much lower than that of the previously reported b-type titanium alloys [3,4,41,42] and similar to that of cortical bone of humans (7e 35 GPa) [25]. Thus, the thermo-mechanical treated Ti19Nb14Zr (at.%) has a great potential for biomedical applications. However, the work hardening would induce defect and/or dislocation in the alloy, the stability of superelasticity during cyclic deformation is important. Hence, further investigation for the stability of superelastic and the fatigue behaviors of this alloy will be studied. 4. Conclusions The effect of thermo-mechanical treatment on shape memory effect and superelastic behavior of Tie19Nbe14Zr (at. %) shape memory alloy were investigated. In both the annealed and aged specimens, a pair of reverse and forward martensitic transformations were detected in their DSC curves. It is found that the formation of small amount of uath phase and/or nanodomains could suppress a large amount of transformations from b phase to a00 phase in the annealed and aged-10 min specimens. Hence, b phase was found to be the predominant phase at room temperature. The volume fraction of u phase increased with increasing the aging time. Interestingly, the volume fraction of both a00 and u phases were simultaneously increased when the specimen further aged for 20 h. It is believed that there are two transformations, b / uiso and uath / a00 , occurred in longer aging time. The annealed specimen exhibited complete shape memory effect up to the maximum strain of w8% after 5 cycles loadinge
unloading test. Non-martensitic superelasticity with approximately of 6% was obtained in the aged-10 min specimen, suggesting that this non-martensitic superelasticity can be obtained in the combination of u þ b phases with extra appropriate precipitation hardening and/or work hardening. Furthermore, the elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately 14 GPa, which is similar to that of cortical bone of human beings (7e35 GPa), thus the thermo-mechanical treated Ti19Nb14Zr (at.%) has a great potential for biomedical applications. Acknowledgment This project is financially supported by the CityU SRG project #7002691. References [1] He G, Hagiwara M. Ti alloy design strategy for biomedical applications. Mater Sci Eng C 2006;26:14e9. [2] Gross S, Abel EW. A finite element analysis of hollow stemmed hip prostheses as a means of reducing stress shielding of the femur. J Biomech 2001;34:995e 1003. [3] Long M, Rack HJ. Titanium alloys in total joint replacementda materials science perspective. Biomaterials 1998;19:1621e39. [4] Niinomi M. Recent metallic materials for biomedical applications. Metall Mater Trans A 2002;33:477e86. [5] Oshida Y, Miyazaki S. Corrosion and biocompatibility of shape memory alloys. Corros Eng 1991;40:1009e25. [6] Miyasaki S, Otsuka K. Development of shape memory alloys. ISIJ Int 1989;49: 353e77. [7] Shabalovskaya SA. Physicochemical and biological aspects of nitinol as a biomaterial. Int Mater Rev 2001;46:233e50. [8] Shabalovskaya S, Cunnick J, Anderegg J, Harmon B, Sachdeva B. Preliminary data on in vitro study of proliferative rat spleen cell response to NieTi surfaces characterized using ESCA analysis. In: Proceedings of the first international conference on shape memory and superelastic technologies; 1994. p. 209. [9] Baker C. The shape-memory effect in a titanium-35 wt.% niobium alloy. Metal Sci J 1971;5:92e100. [10] Kim HY, Hashimoto S, Kim JI, Hosoda H, Miyazaki S. Mechanical properties and shape memory behavior of TieMoeGa alloys. Mater Trans 2004;45: 2443e8. [11] Nitta K, Watanabe S, Masahashi N, Hosoda H, Hanada S. Ni-free TieNbeSn shape memory alloys. In: Proc. Intl. Symp. Structural biomaterials for the 21st century; 2001, p. 25e34. [12] Kim JI, Kim HY, Inamura T, Hosoda H, Miyazaki S. Shape memory characterization of Ti-22Nb-(2-8)Zr(at.%) biomedical alloys. Mater Sci Eng A 2005;403: 334e9. [13] Tahara M, Kim HY, Hosoda H, Nam T, Miyazaki S. Effect of nitrogen addition and annealing temperature on superelastic properties of TieNbeZreTa alloys. Mater Sci Eng A 2010;527:6844e52. [14] Hickman BS. Omega phase precipitation in alloys of titanium with transition metals. Trans Metall Soc AIME 1969;245:1329e36. [15] Hickman BS. The formation of omega phase in titanium and zirconium alloys: a review. J Mater Sci 1969;4:554e63. [16] Bowen AW. Omega phase embrittlement in aged Ti-15%Mo. Scripta Metall 1971;5:709e16.
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Effect of thermo-mechanical treatment on superelastic behavior of Tie19Nbe14Zr (at.%) shape memory alloy L.W. Ma, H.S. Cheng, C.Y. Chung* Department of Physics and Materials Science, City University of Hong Kong, 83 Tat Chee Avenue, Kowloon Tong, Hong Kong
a r t i c l e i n f o
a b s t r a c t
Article history: Received 16 February 2012 Received in revised form 29 May 2012 Accepted 19 July 2012 Available online 9 October 2012
The effect of thermo-mechanical treatment on shape memory effect and superelastic behavior of Ti e19Nbe14Zr (at.%) shape memory alloy is investigated. The annealed specimens are aged at 300 C for different times from 10 min to 20 h followed by ice water quenching. Both reverse and forward martensitic transformation peaks are identified in the annealed and aged specimens. Fine u particles with a dimension of w20 nm in b grains are clearly observed in the TEM micrograph. The strains of the annealed specimen are completely recovered by heating, starting from the 2nd cycle to the 5th cycle due to the reverse martensitic transformation. Non-martensitic superelasticity strain of about 6% is obtained in the aged-10 min specimen. It is noted that the u phase hinders the formation of stress-induced martensite from the b phase, suggesting that the non-martensitic superelasticity should be attributed to the presence of u þ b phases. Also, the non-martensitic superelasticity can be improved through increasing the critical stress for dislocation slip deformation by controlling the size and amount of u phase and/or introducing plastic strain. It is found that this non-martensitic superelasticity can be obtained from the present alloy with combined u þ b phases by introducing additional precipitation hardening and/or work hardening. The elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately of 14 GPa, which is close to that of cortical bone of human beings (7 e35 GPa). Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: B. Martensitic transformations B. Shape-memory effect C. Thermomechanical treatment F. Mechanical testing
1. Introduction Titanium alloys have been widely used for biomedical implants because of their superior corrosion resistance, mechanical properties and tissue acceptance characteristics. However, many studies have reported that unsatisfactory loads transfer from the implant devices and the relatively high elastic modulus of implant materials may lead to bone resorption [1]. The elastic moduli mismatch between the metallic implant and adjacent bone is often referred to the cause of this ‘stress shielding effect’ [2]. Developing alloys with low modulus and superelasticity may provide solutions to ease the stress shielding problem [3,4]. TieNi alloys have been applied to various medical products due to their excellent superelasticity under room and body temperature [5e7]. However, it has been reported that pure Ni is toxic and may cause Nihypersensitivity [8]. Hence, new alloys which are Ni-free have been developed for biomedical applications to substitute TieNi alloys. TieNb based shape memory alloy is one of the candidates which have recently attracted considerable attention for their potential to * Corresponding author. Tel.: þ852 3442 7835; fax: þ852 3442 0538. E-mail address: [email protected] (C.Y. Chung). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.07.024
be used in biomedical devices. It is because they contain no toxic elements and possess superior mechanical properties [9e11]. Many researches have investigated the superelasticity of TieNb based alloys [12,13]. However, the reported superelasticity characteristic is not as good as that of TiNi. Moreover, the incomplete superelastic behavior was unsatisfactory. It is well known that the shape memory effect and superelasticity is based on the reversible b / a00 transformation at loading temperature higher than Af, superelasticity is limited to alloys with Af lower than the temperature of the loading temperature. The transformation peak corresponding to martensitic transformation in TieNb was difficult to be identified using DSC due to small enthalpy of transformation and the large difference between Ms and Mf [12]. Hence, the study and application of Ni-free SMA in human body is not straightforward. Omega (u) phase is a metastable phase in Ti-based alloys with hexagonal structure. It can be easily formed either by quenching from high temperature b phase (athermal-uath) or by aging at intermediate temperatures (isothermal-uiso) [14,15]. It had been reported that the precipitation of u phase had caused embrittlement [16,17]. On the other hand, some papers had suggested that fine u precipitates are effective in improving the superelastic
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property of the titanium alloy [18,19]. It is quite obvious that the presence of u phase in shape memory alloy would affect the shape memory effect and/or superelasticity of titanium alloys; however, it is not well understood. The objective of the present work is to specifically study the effects of thermo-mechanical treatment on superelastic behavior of Tie19Nbe14Zr (at.%) shape memory alloy. The transformation temperatures of the alloy subjected to annealing and different aging time were determined using the Differential Scanning Calorimetry (DSC) thermal analysis. X-ray diffraction (XRD) was used to examine the crystallographic structures of the annealed and aged specimens. Besides, transmission electron microscopy (TEM) was used to reveal the presence of u phase in the aged specimen. Finally, the shape memory effect and superelastic behavior of specimens subjected to both annealing and aging were also investigated by cyclic tensile test. 2. Material and methods The alloy ingot with Tie19Nbe14Zr (at.%) was fabricated by a vacuum consumable arc melting furnace. The ingot was coldrolled to a plate with a reduction in thickness of 90%. Specimens for DSC analysis, XRD measurement, TEM observation and tensile tests were cut from the plate by an electro-discharge machine. The specimens were annealed at 600 C for 30 min followed by ice water quenching. Some of the annealed specimens were then aged at 300 C from 10 min to 20 h followed by ice water quenching. Hereinafter, annealed specimens followed by aging at 300 C with different time-x are abbreviated to aged-x specimens, where x ranged from 10 min to 20 h. An aqueous solution of HF and HNO3 was used to remove the oxidized layer of both the annealed and aged specimens. Phase transformation temperatures were determined between 100 C and 600 C using a Perkin Elmer (PE) Differential Scanning Calorimetry (DSC7) thermal analysis. The specimen weight for DSC thermal analysis was w50 mg, while the heating and cooling rate was 20 K/min. The transformation start and finish temperatures were determined by the intersection of a base line and the tangent to a peak in the DSC curves. The absorbed/released heats during the transformations were calculated from the areas under curves of the heat flows between the start and finish temperatures of transformations. The constituent phases of the specimens were identified using an X’pert MPD Pro XA) ray diffraction (XRD) instrument with Cu-Ka target (l ¼ 1.5406 in 2q ranges between 30 and 80 . TEM specimens were prepared by electrochemical polishing using an electrolyte of 6% perchloric acid þ 35% butanol þ 59% methanol at 10 C. TEM observations were performed using a Philips CM20 microscope. Tensile cycling test and training were carried out at a strain rate of 1.9 104 s1 at room temperature. The gage length of the specimens for tensile tests was 14 mm. Finally, the stress-induced martensite variants of aged specimen after cyclic tensile test were revealed using a JEOL JSM-820 Scanning Electron Microscope (SEM). 3. Results and discussion The transformation profiles of the specimens subjected to different heat treatments were measured by the DSC analysis. All specimens were initially heated to w500 C and then cooled down to w100 C. Fig. 1 shows the DSC curves of (a) the annealed specimen, (b) the aged-1 h specimen and (c) the aged-20 h specimen. All heating and cooling curves were represented using solid lines and dash lines respectively. Two distinct peaks were detected in all specimens. These two peaks are proved a pair of transformation using the “partial DSC cycle technique”, which was usually performed to identify the nature of multi-stage transformation [20]. It
Fig. 1. DSC curves of (a) the annealed specimen, (b) the aged-1 h specimen and (c) the aged-20 h specimens.
is believed that a pair of peaks in each specimen represents the reverse and forward martensitic transformations. Further clarification will be explained by shape memory recovery using stresse strain curves. A slight shift in peak of reverse martensitic transformation was observed in the aged specimens. It is found that both As and Af increase with increase in aging time. It is also noted that the absorbed energy (Eabs) during reverse martensitic transformation increases with longer aging time. The increase in Af and Eabs is possibly due to increasing the stability of martensite by martensitic aging effect [21,22] and/or precipitation of u phase. In contrast, the Ms decreased very slightly with longer aging time. These changes in Ms may be attributed to the precipitation of u phase and/or parent phase aging [23]. Isothermal omega phase (uiso) was reported to have precipitated in b-Ti based alloys when subjected to subsequent aging effect [24]. Subsequently, the formation of uiso may lead to an increase in the b-stabilizer in the matrix, causing the Ms to decrease. The transformation temperatures and energies of the annealed, aged-1 h and aged-20 h specimens obtained from the DSC curves are listed in Table 1. It should be pointed out that the martensitic finish temperatures of all samples are above room temperature. This means that the present alloys should be in martensite state at room temperature. Fig. 2 shows the XRD patterns of the (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen, and (d) aged-20 h specimen. In Fig. 2(a), b phase was surprisingly found to be the predominant phase with the presence of a small amount of u and a00 phases (annealed specimen). Since the Ms and Mf of annealed specimen was measured to be above room temperature (Table 1), expecting that the annealed specimen followed by quenching comprises primarily of a00 martensite rather than b phase. However, only a very weak peak corresponding to a00 martensite was detected. It is
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L.W. Ma et al. / Intermetallics 32 (2013) 44e50
Table 1 Transformation temperatures and energies of the annealed, aged-1 h and aged-20 h specimens.
Annealed specimen Aged-1 h specimen Aged-20 h specimen
Mf
Mp
Ms
Erel
Af
Ap
As
Eabs
138 156 146
212 209 192
250 246 242
2.2 1.8 1.6
419 428 438
384 400 409
310 331 336
14.9 22.6 30.5
Where Mf ¼ martensite finish temperature ( C), Mp ¼ martensite peak temperature ( C), Ms ¼ martensite start temperature ( C), Erel ¼ released energy (J/g), Af ¼ austenite finish temperature ( C), Ap ¼ austenite peak temperature ( C), As ¼ martensite start temperature ( C), Eabs ¼ absorbed energy (J/g).
suggested that the formation of small amount of uath phase could suppress a large amount of transformations from b to a00 , although forward martensitic transformation was observed above room temperature in the DSC curve (Fig. 1(a)). The present observations are similar to those reported by Zain et al. [25] and Lin et al. [26], who investigated the superelastic behavior of TieNbeMo alloys, and the structure and properties of Ti-7.5Mo-xFe alloys respectively. Both groups also found that the formation of small amount of uath phase after quenching suppressed the martensitic transformation. Also, a second stage martensitic transformation [27] may be another possible explanation for the existence of b phase below Mf temperature as a major phase component. For further confirmation, the annealed specimen was initially immersed by liquid nitrogen and then characterized by the XRD analysis. However, the XRD results showed no significant difference with that in Fig. 2(a), the remaining b phase maintained ‘frozen’ or untransformed even at a temperature of about w200 C (liquid nitrogen). Nevertheless, similar ‘frozen’ b structure has been reported in slightly doped Ti50xNi50þx [28] and Ti50(Pd50xCrx) [29] alloys. They both claimed that introducing excess dopants to a shape memory alloy results in a random distribution of a small amount of quasistatic local strain domains or nanodomains which decreases the thermodynamic driving force for forming martensite. The growth process of martensite domains can also be retarded by the pre-existing quasistatic nanodomains. Therefore, b phase was found to be stable at room temperature. In Fig. 2(b), similar XRD spectrum was obtained in the aged10 min specimen as compared with the annealed specimen. However, the volume fraction of u phase increased with decreased in peak intensity of b phase in the aged-1 h specimen (Fig. 2(c)). It is possibly due to the transformation from b to uiso during aging, since the u phase was reported to be easily formed by either quenching (athermal) or aging (isothermal). Interestingly, the volume fraction of both a00 and u phases were simultaneously increased when the
Fig. 2. XRD patterns of the (a) annealed specimen, (b) aged-10 min specimen, (c) aged1 h specimen, and (d) aged-20 h specimen.
specimen aged for 20 h (Fig. 2(d)). It is clear that longer aging time can increase the volume fraction of uiso phase [18,19,30]. Therefore, the specimen aged for 20 h consistently has the largest content of uiso phase. Also, the energy of forward martensitic transformation was decreased from 2.2 to 1.6 J/g, suggesting the uiso phase suppressed a part of martensitic transformation by longer aging time. However, the volume fraction of a00 phase also increased with longer aging time should be associated with the transformation from uath back to a00 martensite. Hence, it is suggested that there are two transformations, b / uiso and uath / a00 , occurred in the aged specimen for 20 h. This is in agreement with the phase transformation of a00 martensite structure in aged Tie8wt%Mo alloy reported by Mantani and co-workers [21]. They reported that a00 twins were formed when the u phase products were changed into the a00 martensite by aging. Fig. 3 shows (a) the bright-field TEM image of the aged-1 h specimen, (b) the corresponding selected area diffraction pattern (SADP) of Fig. 3(a) and (c) the key diagram of diffraction pattern. Fine u particles with a dimension of w20 nm in b grains were clearly observed. The selected diffraction pattern was obtained from the [110]b zone axis. The diffraction pattern of the b þ u matrix is consistent with the work of Kim et al. [18]. Strong u reflections in the diffraction pattern indicated that there is a high volume fraction of u phase in the selected area. Also, extra reflections resulting from double diffraction [31] were observed in the pattern, it has been reported in previous studies that the presence of u phase in b grains was always accompanied with broadening of diffraction spots and diffuse streaks [32e34] resulting from the incomplete (111) plane collapse from b phase. Tensile tests were performed at room temperature to study the effect of thermo-mechanical process to shape memory properties of the specimens. Fig. 4 shows the stressestrain curves of (a) asrolled specimen, (b) annealed specimen (c) aged-10 min specimen and (d) aged-1 h specimen. The as-rolled specimen exhibit a single stage yielding with only small plastic deformation and high critical stress compared with other heat-treated specimens. It is due to the work hardening effect induced by cold rolling. Both of the annealed (Fig. 4(b)) and aged10 min (Fig. 4(c)) specimens exhibit a two-stage yielding [12]. It should be noted again that b phase was observed as a predominant phase in all specimens according to the XRD patterns shown in Fig. 2. Typically, stress induced martensitic transformation occurs at a temperature above Af in b phase. However, the b phase is ‘frozen’ even at a temperature below Mf (i.e. room temperature), it is believed stress induced martensitic transformation also occurs in this frozen b structure at room temperature [28,29]. Hence, the apparent yielding stress (first stage yielding) mainly corresponds to the critical stress for inducing the martensitic transformation upon loading (sSIM). Since the As of the alloy is much higher than the loading temperature (i.e. room temperature) as detected by DSC, the stress-induced martensites were maintained. The slope after the apparent yielding stress is related to the transformation of stress-induced martensite (MSI) to martensite variants (Mv) (second stage yielding). In Fig. 4(c), the slope for the transformations of MSI / Mv and critical stress for slip increased after aging as compared with the annealed specimen (Fig. 4(b)). The ductility of the aged-10 min specimen was decreased by this hardening effect and it is apparently due to the uiso precipitation as confirmed by the XRD profiles or the structural change of a00 martensite [21]. However, the aged-1 h specimen did not exhibit two-stage yielding, though there is minor difference in the XRD profiles between the annealed and the aged-1 h specimens. Furthermore, the critical stress of the aged-1 h specimen is the highest one among all specimens. It is suggested that the uiso precipitates are effective in increasing the critical stress for forming
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Fig. 3. (a) The bright-field TEM image of the aged-1 h specimen, (b) the corresponding selected area diffraction pattern (SADP) of Fig. 3(a) and (c) the key diagram of diffraction pattern.
stress-induced martensite and the stress for the transformation of M / Mv. Therefore, no significant amount of stress-induced martensite and martensite variants were formed before plastic deformation. It is also found that the elastic modulus of the specimen is sensitive to the aging time. The elastic modulus of the annealed specimen increases from 14 GPa to 23 GPa after aging for 1 h. This indicates that the elastic modulus increases with
Fig. 4. Shows the stressestrain curves of (a) as-rolled specimen, (b) annealed specimen (c) aged-10 min specimen and (d) aged-1 h specimen.
increasing the volume fraction of u phase as the Young’s modulus of u phase is much higher than that of b and a00 phase [35]. To further investigate the shape memory and superelastic properties of the alloy, cyclic loading and unloading tensile test was performed at the room temperature. Fig. 5 shows the stressestrain curves obtained by cyclic loadingeunloading tensile test of (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen and (d) the aged-10 min specimen after five cycles of loadingeunloading tests in Fig. 5(c). The annealed specimen was loaded up to 2% in strain and then unloaded during the first cycle, Fig. 5(a). The measurement was repeated by increasing the maximum strain by 1% upon loading for the same annealed specimen followed by heating to a temperature higher than Af, i.e. 450 C, after each cycle. The shape memory recovery and superelasticity were confirmed in the annealed specimen, a line with an arrow indicates the shape recovery by heating. The shape memory recovery of the annealed specimen (Fig. 5(a)) occurred only when the specimen reached a temperature of about 450 C which is consistent with the peak on the heating curve of DSC result (Fig. 1(a)), hence it is believed that the peak on the heating curve is due to reverse martensitic transformation while the peak obtained during cooling corresponds to the forward martensitic transformation. The strains are completely recovered by heating, starting from the 2nd cycle to the 5th cycle due to the reverse martensitic transformation. The strains obtained in the curves can be defined into three types: (1) the elastic strain (3 el) recovered
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Fig. 5. The stressestrain curves obtained by cyclic loadingeunloading tensile test for (a) annealed specimen, (b) aged-10 min specimen, (c) aged-1 h specimen and (d) the aged10 min specimen which already subjecting to five cycles of loadingeunloading tests in Fig. 5(b).
elastically upon unloading, (2) the strain recovered superelastically upon unloading (3 se) and (3) the strain recovered by heating (3 sme). After five cycles, the annealed specimen exhibits complete shape memory effect and superelasticity up to the maximum strain of w8%. Hence the sum of 3 se, 3 el, 3 sme as shown in Fig. 5(a) can be indicated as the total recoverable strain (3 r). It should be noted that the superelasticity of SMA is usually related to stress-induced reversible martensitic transformation. In the present study, the superelasticity (non-martensitic) obtained was apparently not associated with the stress-induced reversible martensitic transformation since the reverse martensitic transformation temperatures of the present alloy were confirmed to be much higher than room temperature by DSC analysis (Fig.1). It was reported that “nonmartensitic” alloy Ni48.5Ti51.5 can exhibit both SME and SE without showing martensitic transformation [36]. Furthermore, nonmartensitic superelasticity can also be resulted from other kinds of mechanisms, such as dislocation based mechanism of incipient kink band (IKB) [37] and dislocation free mechanism of giant fault [38] or dislocation-free sliding. In Fig. 5(b), almost complete nonmartensitic superelasticity was observed in aged-10 min specimen during the first and second cycles. However, only single stage yielding was obtained in aged-1 h specimen (Fig. 5(c)). It is suggested that the critical stress for inducing martensite is much higher than the critical stress for slip deformation. A severe loss of ductility was obtained in the aged-1 hr. specimen compared with the aged-10 min specimen. This implies that the precipitations of more thermal u phase formed after longer aging time is effective in hardening. Hence, it is considered that Tie 19Nbe14Zr is not a promising high temperature shape memory alloy as it leads to brittleness when subjecting at high temperature. In contrast, an extra cyclic loading and unloading test (Fig. 5(d)) was specifically performed on the same aged-10 min specimen which already subjecting to five cycles of loadingeunloading test in Fig. 5(c). Complete non-martensitic superelasticity with approximately of 6% was obtained.
In this investigation, the actual mechanism for the “nonmartensitic superelasticity” obtained in the present alloy should be associated with the combination of u þ b phases and/or the atomic rearrangement within the same sub-lattice of the imperfectly ordered alloy during martensite aging [22]. It is suggested that the u phase hinders the formation of stress-induced martensite from the b phase. Hence, the non-martensitic superelasticity can be improved in this combination of u þ b phases with extra appropriate precipitation hardening (aging) and/or strain hardening. The critical stress for slip deformation (sCSS), the critical stress for inducing martensite (sSIM) and the elastic modulus (E) for aged10 min specimen estimated in Fig. 5(b) was plotted against the maximum strain (%) in Fig. 6.
Fig. 6. The critical stress for slip deformation (sCSS), the critical stress for inducing martensite (sSIM) and the elastic modulus (E) for aged-10 min specimen estimated in Fig. 5(b) was plotted against the maximum strain (%).
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Fig. 7. SEM micrographs of the aged-10 min specimen (a) before the cyclic tensile test, and (b) after the cyclic test with maximum strain of 8%.
The critical stress for slip deformation (sCSS) increased continuously, should be associated with the work hardening effect during mechanical cycling. However, the critical stress for inducing martensite (sSIM) decreased with increasing the maximum loading strain. The elastic modulus is approximately of 14 GPa and decreases slightly with increasing maximum strains. The increase in volume fraction of residual a00 phase can be attributed to stress induced martensitic transformation upon loading and unloading. SEM examination was used to study the structural change of the alloy after the cyclic tensile test. Fig. 7 shows the SEM micrographs of the aged-10 min specimen (a) before the cyclic tensile test, and (b) after the cyclic test with maximum strain of 8%. Stress-induced martensite variants with lamellar structures were observed within the equi-axed b grains. It is clear that there is a transformation of almost entire b phase to a b þ a00 system due to the formation of stress induced martensite. Also, the elastic modulus of the aged-10 min specimen decreased slightly upon cyclic loading due to the lower modulus nature of the a00 phase [39,40]. The elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately of 14 GPa, which is much lower than that of the previously reported b-type titanium alloys [3,4,41,42] and similar to that of cortical bone of humans (7e 35 GPa) [25]. Thus, the thermo-mechanical treated Ti19Nb14Zr (at.%) has a great potential for biomedical applications. However, the work hardening would induce defect and/or dislocation in the alloy, the stability of superelasticity during cyclic deformation is important. Hence, further investigation for the stability of superelastic and the fatigue behaviors of this alloy will be studied. 4. Conclusions The effect of thermo-mechanical treatment on shape memory effect and superelastic behavior of Tie19Nbe14Zr (at. %) shape memory alloy were investigated. In both the annealed and aged specimens, a pair of reverse and forward martensitic transformations were detected in their DSC curves. It is found that the formation of small amount of uath phase and/or nanodomains could suppress a large amount of transformations from b phase to a00 phase in the annealed and aged-10 min specimens. Hence, b phase was found to be the predominant phase at room temperature. The volume fraction of u phase increased with increasing the aging time. Interestingly, the volume fraction of both a00 and u phases were simultaneously increased when the specimen further aged for 20 h. It is believed that there are two transformations, b / uiso and uath / a00 , occurred in longer aging time. The annealed specimen exhibited complete shape memory effect up to the maximum strain of w8% after 5 cycles loadinge
unloading test. Non-martensitic superelasticity with approximately of 6% was obtained in the aged-10 min specimen, suggesting that this non-martensitic superelasticity can be obtained in the combination of u þ b phases with extra appropriate precipitation hardening and/or work hardening. Furthermore, the elastic modulus of the aged-10 min specimen after aging and cyclic training is approximately 14 GPa, which is similar to that of cortical bone of human beings (7e35 GPa), thus the thermo-mechanical treated Ti19Nb14Zr (at.%) has a great potential for biomedical applications. Acknowledgment This project is financially supported by the CityU SRG project #7002691. References [1] He G, Hagiwara M. Ti alloy design strategy for biomedical applications. Mater Sci Eng C 2006;26:14e9. [2] Gross S, Abel EW. A finite element analysis of hollow stemmed hip prostheses as a means of reducing stress shielding of the femur. J Biomech 2001;34:995e 1003. [3] Long M, Rack HJ. Titanium alloys in total joint replacementda materials science perspective. Biomaterials 1998;19:1621e39. [4] Niinomi M. Recent metallic materials for biomedical applications. Metall Mater Trans A 2002;33:477e86. [5] Oshida Y, Miyazaki S. Corrosion and biocompatibility of shape memory alloys. Corros Eng 1991;40:1009e25. [6] Miyasaki S, Otsuka K. Development of shape memory alloys. ISIJ Int 1989;49: 353e77. [7] Shabalovskaya SA. Physicochemical and biological aspects of nitinol as a biomaterial. Int Mater Rev 2001;46:233e50. [8] Shabalovskaya S, Cunnick J, Anderegg J, Harmon B, Sachdeva B. Preliminary data on in vitro study of proliferative rat spleen cell response to NieTi surfaces characterized using ESCA analysis. In: Proceedings of the first international conference on shape memory and superelastic technologies; 1994. p. 209. [9] Baker C. The shape-memory effect in a titanium-35 wt.% niobium alloy. Metal Sci J 1971;5:92e100. [10] Kim HY, Hashimoto S, Kim JI, Hosoda H, Miyazaki S. Mechanical properties and shape memory behavior of TieMoeGa alloys. Mater Trans 2004;45: 2443e8. [11] Nitta K, Watanabe S, Masahashi N, Hosoda H, Hanada S. Ni-free TieNbeSn shape memory alloys. In: Proc. Intl. Symp. Structural biomaterials for the 21st century; 2001, p. 25e34. [12] Kim JI, Kim HY, Inamura T, Hosoda H, Miyazaki S. Shape memory characterization of Ti-22Nb-(2-8)Zr(at.%) biomedical alloys. Mater Sci Eng A 2005;403: 334e9. [13] Tahara M, Kim HY, Hosoda H, Nam T, Miyazaki S. Effect of nitrogen addition and annealing temperature on superelastic properties of TieNbeZreTa alloys. Mater Sci Eng A 2010;527:6844e52. [14] Hickman BS. Omega phase precipitation in alloys of titanium with transition metals. Trans Metall Soc AIME 1969;245:1329e36. [15] Hickman BS. The formation of omega phase in titanium and zirconium alloys: a review. J Mater Sci 1969;4:554e63. [16] Bowen AW. Omega phase embrittlement in aged Ti-15%Mo. Scripta Metall 1971;5:709e16.
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