Phase transformation during aging and resulting mechanical ...

Phase transformation during aging and resulting mechanical properties of two Ti–Nb–Ta–Zr alloys S. J. Li1, R. Yang*1, M. Niinomi2, Y. L. Hao1, Y. Y. Cui1 and Z. X. Guo3 Phase transformations and mechanical properties of both Ti–29Nb–13Ta–4?6Zr and Ti–39Nb– 13Ta–4?6Zr (wt–%) alloys were investigated. The microstructure of the 29Nb alloy is sensitive to solution and aging treatment. Ice water quenching from the solution treatment temperature resulted in (bza’’) microstructure but air or furnace cooling led to a mixture of (bzv). The formation of the orthorhombic a’’ martensite thus suppresses v formation in the ice water quenched 29Nb alloy. Cooling rate from the solution treatment temperature also has a significant effect on the formation of a and v phases during subsequent isothermal aging below the v start temperature: slow cooling enhances v but depresses a formation. This cooling rate dependence of aged microstructure was attributed to a’’ martensite acting as precursor of the a phase, thus providing a low energy path to the precipitation of a at the expense of v. Phase transformation in the 39Nb alloy is more sluggish than that in the 29Nb alloy, owing to the presence of the higher content of b stabiliser Nb. For the 29Nb alloy, Young’s modulus and mechanical properties are sensitive to the fraction of phases, and change significantly during aging, in contrast with the 39Nb alloy. Keywords: Metastable b titanium alloy, phase transformation, Young’s modulus, mechanical properties

Introduction Orthopaedic implant applications of materials require good mechanical properties and biochemical compatibility. Due to the low density, super biocompatibility and biocorrosion resistance, good mechanical properties, and low elastic modulus of titanium alloys, they are generally preferred to stainless steels and Co–Cr alloys and have become one of the most attractive classes of candidate implant materials. For example, Ti–6Al–4V ELI (extra low interstitial), having an elastic modulus only about half that of 316 stainless steel or Co–Cr–Mo alloy, has been widely used as implant material. Recent concerns over potential long term health problems caused by the release of Al and V ions from the alloy, and on load shedding due to differences of stiffness between bone and implant1,2 have led to an increased level of interest in novel titanium alloys with better biocompatibility, still lower modulus, and better 1

Titanium Alloy Laboratory, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China Department of Production Systems Engineering, Toyohashi University of Technology, 1–1, Hibarigaoka, Tempaku–cho, Toyohashi 441–8580, Japan 3 Department of Materials, Queen Mary, University of London, Mile End Road, London E1 4NS, UK 2

*Corresponding author, email [email protected]

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processability. As a result, near b or metastable b type Ti–Nb based titanium alloys with Zr and/or Ta additions have been developed.3–5 Moffat et al. conducted an extensive investigation of phase transformations in the binary Ti–Nb system.6–8 Their results showed that the formation of a0 and v phases varies with both quench rate and chemical composition during quenching after solution treatment. The a0 martensite was found below 39% Nb (wt-%, as are all chemical compositions quoted below) under the condition of fast water quenching, whereas trigonal vath and ideal hexagonal v phases can form in alloys with Nb ranging between 39% and 66%.6 During aging treatment, stable a phase and metastable v phase can be formed in Ti–Nb alloys, depending on chemical composition and aging temperature.7 Ahmed and Rack 9 also focused on the formation of a0 martensite in Ti–(28–40)Nb alloys and they reported an ordered base centred orthorhombic a0 martensite at Nb contents up to 37?2%. Their results suggested that both the martensite start (Ms) temperature and the order– disorder temperature are sensitive to the contents of Nb and interstitials, and decrease with increasing amount of Nb and OzNzC. Phase transformations in Ti–Nb–Ta and Ti–Nb–Ta– Zr alloys have been investigated very recently. The experimental results of Tang et al.10 showed that phase

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transformations are sensitive to both alloy composition and cooling rate. Orthorhombic a0 martensite forms in water and oil quenched specimens but fine a and v form at lower cooling rate. The increase of NbzTa content decreases the volume percentage of martensite, while Zr tends to decrease martensite start temperature and suppress v formation. Hao et al.11,12 evaluated phase transformations in forged Ti–29Nb–13Ta–4.6Zr alloy based on X-ray and internal friction measurements. The results suggested that, for ice water quenched specimens after solution treatment, the formation of a0 martensite depends on both solution temperature and time. Systematic investigation of the microstructure of Ti– 29Nb–13Ta–4?6Zr alloy by means of optical and electron microscopy analysis, however, is lacking, particularly in the aged conditions. In the present study, we investigate phase transformations of both Ti–29Nb–13Ta–4.6Zr and Ti–39Nb– 13Ta–4.6Zr alloys and evaluate the effect of Nb on phase transformations. The influence of cooling rate after solution treatment and aging temperature on the phase transformations, Young’s modulus, and mechanical properties of Ti–29Nb–13Ta–4.6Zr alloy were also examined. The relationship among the different phases is paid particular attention in an attempt to elucidate the sequence of phase transformations that produces the observed microstructures.

Experimental A Ti–29Nb–13Ta–4?6Zr ingot with diameter 60 mm was fabricated by induction skull melting using pure Ti, Nb, Ta and Zr as raw materials and then hot forged to rods with a diameter of 20 mm. A Ti–39Nb–13Ta–4?6Zr ingot was produced by first melting in an electron beam melter using raw materials of pure Ti, Nb, Ta, and Zr, followed by remelting in a vacuum arc furnace. The ingot was then forged at 850uC to a rod 24 mm in diameter. The compositions of the two alloys in the asforged state according to chemical and gas analysis are shown in Table 1. Cylindrical sections 25 mm in length were taken from the two alloy rods for microstructure analysis. The specimens for test of Young’s modulus and mechanical properties were taken from rectangular plates 6261263 mm in size and those used for phase identification analysis by X-ray diffraction were 1061063 mm plates sliced along cross and longitudinal directions. The heat treatments of all specimens were carried out under argon protection. The specimens were first solution treated at 790uC for 1 h followed by ice water quenching (WQ). The solution temperature is higher than the b transus of the two studied alloys: 650uC for Ti–29Nb–13Ta–4?6Zr estimated by X-ray and internal friction measurements,12 and 550uC for Ti– 39Nb–13Ta–4?6Zr estimated by considering alloying effects.13 After WQ treatment, the specimens were aged Table 1 Chemical compositions of forged alloys Composition, wt-% Alloy

Nb

Ta

Zr

O

N

Ti

Ti–29Nb–13Ta–4.6Zr Ti–39Nb–13Ta–4.6Zr

29.2 39.8

12.2 13.5

4.3 5.0

0.10 0.088

0.04 bal. 0.015 bal.

Mechanical properties of two Ti–Nb–Ta–Zr alloys after aging

at 300, 350, 400, 450, or 500uC for 48 h and then quenched in ice water. To investigate the effect of cooling rate on microstructure and mechanical properties, several Ti–29Nb–13Ta–4?6Zr specimens were also air cooled (AC) or furnace cooled (FC) after solution treatment and then aged at 300, 400, or 500uC for 48 h. Based on the data reported by Moffat et al.,7 cooling rate from 790uC can be estimated as 300, 5, and 0?3 K s–1 for WQ, AC, and FC, respectively. Rectangular plates of 6061062 mm dimension were machined and ground after heat treatments, and then used to measure the dynamic Young’s modulus by the resonance method at ambient temperature. The Young’s modulus E was calculated using the expression14 E~0:9694rL4 f 2r =d 2 where r, L, and d are the density, length and thickness of specimens, respectively, and fr is resonance frequency. The length and thickness of specimens were measured using an Olympus microscope. The error in the dynamic Young’s modulus is approximately ¡2?5% of its actual value, which was estimated from the errors in specimen dimension caused by grinding and polishing.11 Specimens of 2 mm thickness, 3 mm width and 12 mm gauge length were used to measure the tensile properties at an initial strain rate of 761024 s21. Vickers hardness was measured with a load of 10 kg applied for 15 s. The microstructures were observed with a JSM– 6301F scanning electron microscope (SEM) and a Philips EM420 transmission electron microscope (TEM) operating at 100 kV. The specimens for SEM analysis were mechanically polished and then etched in a solution consisting of 8 vol.-%HF, 15 vol.-%HNO3 and 77 vol.-%H2O. Samples for TEM analysis were prepared from mechanically thinned plates by ion milling with 5 keV argon ions at an angle of incidence of 15u without cold finger attachment.

Results Microstructures of as solution treated alloys Figure 1a shows an electron micrograph of the (bza0) microstructure of Ti–29Nb–13Ta–4?6Zr alloy solution treated at 790uC for 1 h followed by WQ. X-ray diffraction analysis confirmed the presence of (020), (111) and (113) peaks of the orthorhombic a0 martensite in WQ specimens (Fig. 1b). Details of the a0 martensite formation can be found in Ref. 11. For WQ specimens of Ti–39Nb–13Ta–4?6Zr, no evidence of orthorhombic a0 martensite was found by transmission electron microscopy, due to increased stability of the b phase and lower martensite start (MS) temperature as a result of the increased content of b stabiliser Nb, in agreement with previous experimental results.6,10 Diffuse scattering was observed on selected area electron diffraction (SAD) patterns in both alloys under WQ condition (Fig. 2a and b). Compared with the high Nb alloy (Fig. 2a), the low Nb alloy exhibits weaker diffuse streaking and the diffuse maxima are not yet formed (Fig. 2b). The diffuse maxima in the high Nb alloy are identified to be due to the trigonal v structure as observed in the Ti–Nb15 and TiAl–Nb alloys.16 Therefore, the most likely origin of the diffuse scattering is precursor of the athermal transformation to the v


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the v phase can be determined as [113]b|| [12¯13]v1|| [1¯1¯23¯]v2. X-ray analysis also showed that the a phase, which has needle like morphology (Fig. 6), was formed when aging was conducted above 350uC for 2 days (Fig. 4). The fine size of the a and v particles prevented their differentiation by optical metallographic analysis, so an X-ray diffraction technique was used to study the change of their volume fractions with aging temperature in Ti–29Nb–13Ta–4?6Zr. The method described by Lopata and Kula17 was used, with which the relative amounts of phases were calculated according to IA =IB ~(RA =RB )(CA =CB )

a bright field TEM micrograph; b X-ray diffraction pattern (sample taken in the transverse orientation) 1 Microstructure and phase identification of water quenched Ti–29Nb–13Ta–4?6Zr

structure. Figure 2a also shows that the diffuse maxima are shifted away from the ideal positions of the v reflections, indicating the embryo nature of the v type phase. Direct observation failed to reveal the v phase particles in both alloys. SAD patterns of air and furnace cooled Ti–29Nb– 13Ta–4?6Zr specimens are given in Fig. 2c and d, respectively. Comparing with those of WQ specimens (Fig. 2a and b), diffuse maxima are stronger and closer to the ideal v positions with decreased cooling rate. For AC specimens, a dark field image of the v phase is not yet obtainable. Furnace cooling resulted in sharper and stronger v phase reflections on the SAD patterns (Fig. 2d) and v particles can be definitely observed on the dark field micrograph (Fig. 3), suggesting that ideal hexganal v phase formed in FC samples.

Microstructures of aged Ti–29Nb–13Ta–4?6Zr Effect of aging temperature. The microstructure of aged Ti–29Nb–13Ta–4?6Zr is very sensitive to aging temperature. The v phase was formed in specimens aged below 400uC for 2 days, evidenced by X-ray diffraction patterns presented in Fig. 4. Dark field TEM shows that the v particles have ellipsoidal shape and distribute uniformly in the b grains (Fig. 5a). Based on the SAD pattern obtained from [113]b zone axis given in Fig. 5b, the orientation relationship between the bcc matrix and

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where Ii is the integrated intensity of peak i of a, v or b phase, Ci is the concentration of phase i in the mixture, and Ri is a geometrical factor depending on the atomic volume, structure factor, multiplicity factor, temperature factor and a factor depending on the Bragg angle 2h. Previous investigation showed that the forged Ti– 29Nb–13Ta–4?6Zr rod has strong n110m texture along its longitudinal direction and heat treatment only has slight influence on the texture.11 The height ratios of Xray peaks of the a and the v phases with respect to the b phase can, therefore, be used to describe the variation of volume fractions of the a and the v phases with aging temperature. Since X-ray analysis was carried out along longitudinal direction of the rod, the height ratios of Ia(110)/Ib(211) and Iv(001)/Ib(211) were measured and the results are plotted in Fig. 7. It is clear that with the decrease of aging temperature from 500 to 400uC, the volume fraction of the a phase increases until about 450uC and then decreases with further decrease in aging temperature; the volume fraction of the v phase increases from 400 to 300uC whereupon the maximum amount appears to have been reached. If the peaks of the a and the v phases are compared against (110)b, a similar conclusion can be reached, agreeing with that obtained previously based on internal friction measurement.12 For Ti–39Nb–13Ta–4?6Zr alloy with the higher Nb content, aging between 300 and 500uC for 2 days produces no definite a reflections on the SAD patterns, while only very faint and weak v reflections were observed in the samples aged below 400uC.13 The sluggish phase transformations from the b to the a and to the v phase are attributed to the increased stability of the b matrix as a result of increased content of the b stabilising Nb. Additional experiment reveals that recognisable a phase forms when the sample is aged at 500uC for 288 h,13 suggesting that Ti–39Nb–13Ta–4?6Zr also belongs to the metastable b type titanium alloys. Effect of cooling method from solution treatment. After aging at 500uC for 48 h, only a precipitation was detected in WQ, AC, and FC specimens by X-ray diffraction analysis (Fig. 8). In comparison with the microstructure of the WQ and aged sample, the size of the a phase is much finer in the AC and aged sample (Fig 6c and d), whereas the FC and aged sample exhibits coarse a plates (Fig. 6e and f). Aging at 400uC produced an interesting result: While precipitation of both a and v occurred, the amount of v phase increases with decreasing cooling rate and that of a decreases correspondingly. It should be stated that this conclusion

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a Ti–39Nb–13Ta–4?6Zr, WQ; b Ti–29Nb–13Ta–4?6Zr, WQ; c Ti–29Nb–13Ta–4?6Zr, AC; d Ti–29Nb–13Ta–4?6Zr, FC 2 Selected area electron diffraction patterns in the [011]b zone

is not compromised by the presence of crystallographic texture because all diffraction peaks belonging to the same phase in Fig. 8 show a similar trend of variation. In samples aged at 300uC, only v precipitation was detected, the amount of which appears independent of cooling rate. The v particles in these samples are ellipsoidal in shape (similar to that shown in Fig. 5a), although becoming slightly coarser with decrease of cooling rate.

Young’s modulus and mechanical properties Figure 9 presents the Young’s modulus and mechanical properties of the two Ti–Nb–Ta–Zr alloys aged at different temperatures. Significant changes in the

properties of Ti–29Nb–13Ta–4?6Zr are noted, whereas the properties of the high Nb alloy change slightly. This is in accordance with the microstructural features described above, that is, the b matrix is more stable in the high Nb alloy. The Young’s modulus and mechanical properties of aged Ti–29Nb–13Ta–4?6Zr cooled from the solution treatment temperature by different methods are compared in Fig. 10. After aging at 400uC for 2 days, the Young’s modulus and strength (Young’s modulus in particular) increased with decrease of the cooling rate from the solution treatment temperature, whereas the elongation decreased. For aging at 500uC, the AC and aged sample exhibited the highest Young’s modulus and strength, whereas for samples aged at 300uC, the Young’s modulus and mechanical properties did not show obvious change with decrease of cooling rate.

Discussion Formation of v phase

3 Dark field TEM micrograph Ti–29Nb–13Ta–4?6Zr

of

furnace

cooled

Influence of a0 martensite on v formation during cooling. The observation that v like diffuse scattering is stronger in 39Nb alloy than in 29Nb alloy under the condition of ice water quenching (Fig. 2a and b) appears to contradict the argument based on the stability of the b phase (which increases with increasing amount of b stabiliser Nb). Because cooling rate as well as contents of Ta and Zr are identical for both alloys, interpretation of the above observation is probably related to the formation of a0 martensite. Ahmed and Rack9 suggested that the formations of martensite and v phases in binary Ti–Nb alloys are


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a 300uC; b 350uC; c 400uC; d 450uC; e 500uC 4 X-ray diffraction patterns of Ti–29Nb–13Ta–4?6Zr solution treated at 790uC for 1 h, water quenched and then aged for 2 days at different temperatures (samples sectioned in longitudinal orientation)

determined by the martensite start (Ms) temperature and the omega start (Vs) temperature. If the conditions of Vs being above Ms and Ms being above room temperature are both satisfied, v formation is favoured by slow cooling but martensite by rapid cooling. For the low Nb alloy we used, Vs was around 400uC;18 although Ms was not measured, the absence of a0 peaks in the X-ray diffraction spectrum of samples aged at 300uC (Fig. 8) and its occurrence upon WQ to room temperature (Fig. 1) suggests that Ms lies between 300uC and room temperature. Thus the formation of a0 martensite in the low Nb alloy under WQ condition can be understood. While a0 martensite can still form in the presence of v phase particles in Ti–10V–2Fe–3Al as noted by Duerig et al.19, it is clear that in the present alloy the formation of a0 martensite obliterates the v phase. In the high Nb alloy, no martensite was observed under the WQ condition because its Ms is below room temperature, leading to stronger v like diffuse scattering and maxima than in the low Nb alloy. Effect of cooling rate after solution treatment on v formation during aging. Because the formation of athermal v phase involves a displacive transformation by collapse of (111) atomic planes,7 it should not depend on cooling rate. The observed cooling rate dependence of v formation therefore suggests an isothermal character, especially during furnace cooling (Fig. 3). Since the formation mechanism of both isothermal and athermal v phase is identical,20,21 no distinction will be made between the two in the discussion below.

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a dark field TEM micrograph on v reflection; b [113]b SAD pattern; c index of (b) 5 Microstructure of Ti–29Nb–13Ta–4?6Zr water quenched and then aged at 300uC for 2 days

The interesting observation of the varying amount and morphology of precipitated phases in aged samples with different cooling history must be addressed. For Ti–29Nb–13Ta–4?6Zr, v phase forms below about 400uC whereas a phase precipitates above about 350uC. Thermodynamic equilibrium requires that a given phase has identical volume fraction at a given temperature irrespective of the history of the samples. This is indeed the case for aging at 500uC or 300uC, whereupon the volume fraction of the a or v phase, respectively, appears very similar for WQ, AC and FC samples, judging from the strength of the X-ray diffraction peaks shown in Fig. 8. It has been suggested that the athermal v phase would accelerate the precipitation kinetics of the isothermal v phase below the v start temperature.10 The data of the present alloy aged at 300uC (Fig. 8) does not support this assumption. For aging at 400uC, however, the surprising dependence of the volume fractions of v and a phases on the history of specimens, namely the manner of cooling

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a, b WQ; c, d AC; e, f FC 6 SEM micrographs of Ti–29Nb–13Ta–4?6Zr solution treated at 790uC for 1 h, followed by different cooling to room temperature before aging at 500uC for 2 days. Low magnification photographs on the left show general views across b grain boundaries, while high magnification images on the right show details of a precipitates

from solution treatment temperature, may be related to the occurrence of the a0 martensite. The a0 martensite is similar to the orthorhombic c phase observed during high pressure phase transformation of pure titanium,22 and its orthorhombic structure may be viewed as an intermediate step between the bcc and the hcp structure, much like their ordered versions in the intermetallic Ti–Al–Nb system.23 In pure titanium, the rthorhombic c phase is an intermediate phase during the transformation from the v to the b phase with increasing pressure;22 in multicomponent alloys, however, the v phase and a0 martensite are allowed to coexist by thermodynamics, and their formation may be competitive. When the 29Nb alloy is aged at 400uC, the a0 in the WQ samples (Fig. 1) may act as precursor for the equilibrium a phase, overcoming the kinetic difficulties

at such a low temperature. This meanwhile makes the competing process of v formation difficult. In the FC samples, v phase was already copious (Fig. 3) in the as cooled state, and the lack of a0 makes the precipitation of the equilibrium a phase difficult and slow. The AC samples, which show neither a0 martensite nor v particles that can be imaged, allow formation of both phases when aged at 400uC. The a0 martensite, which becomes unstable above Ms and reverts back to the parent b phase when the formation of a is not possible below 350uC, does not influence the formation of the v phase, evidenced by diffraction data collected from samples aged at 300uC (Fig. 8). When aged at 500uC, both a0 and v phases becomes unstable, giving way to the equilibrium a phase. While the volume fraction of the a phase is similar for


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produce the finest a particles during aging at 500uC (Fig. 6c, d)). These variations in microstructure are responsible for the difference in mechanical properties presented in Fig. 10.

Relation of Young’s modulus to phase fractions

7 Variations of amount of a and v phase with aging temperature in Ti–29Nb–13Ta–4?6Zr solution treated (WQ) and then aged for 2 days, estimated by Ia(110)/Ib(211) and Iv(001)/Ib(211) calculations, respectively

samples of three cooling conditions, the differences arising from the cooling history are reflected in the distinct morphology of the a precipitates (Fig. 6) as well as resulting mechanical properties (Fig. 10). Thus the uniform needle like a precipitates of the WQ samples (Fig. 6a, b) derive from the a0 martensite (Fig. 1), whereas the coarse a particles (fewer in number) in the FC samples (Fig. 6e, f)) may have been developed from a limited number of nuclei. As suggested by Boyer and Lu¨tjering,24 v phase formed during cooling after solution treatment may restrict the nucleation of the a particles. Again, the AC samples, which contain neither a0 phase nor obvious v phase before aging was started,

Because the strength and Young’s modulus of an alloy are related to the relative amount of its constituent phases,25–31 it is appropriate to discuss the mechanical properties and Young’s modulus with reference to microstructure of the alloy formed during aging. The aging treatment produces two phases, a and v, from the b matrix of Ti–29Nb–13Ta–4?6Zr. Of these phases, v has the highest Young’s modulus and lowest ductility. Compared to the b phase, the a phase is harder and possesses higher Young’s modulus than the b.31,32 In the aging regime studied, with decrease in aging temperature from 500 to 400uC, the volume percentage of the a phase increased up to 450uC and then decreased (Fig. 7). As a result, Young’s modulus, strength, and hardness (Fig. 9) increased first and then decreased in this temperature range. With further decrease of aging temperature from 400 to 300uC, the volume percentage of v phase increased at the expense of a phase. Accordingly, the strength and hardness of Ti–29Nb–13Ta–4?6Zr decreased due to the loss of a phase, whereas Young’s modulus increased due to the increase in the amount of v phase.

Conclusions 1. Ice water quenching after solution treatment at 790uC for 1 h results in a (bza0) microstructure in

a WQ; b AC; c FC 8 X-ray diffraction patterns of Ti–29Nb–13Ta–4?6Zr solution treated at 790uC for 1 h, followed by different cooling regimes to room temperature, and then aging for 2 days at (1) 300uC; (2) 400uC; (3) 500uC

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Ti–29Nb–13Ta–4?6Zr and (bzv) phases in Ti–39Nb– 13Ta–46Zr. Decrease of cooling rate from solution temperature promotes v precipitation in as solution treated Ti–29Nb–13Ta–4?6Zr. 2. Phase transformations in ice water quenched Ti– 29Nb–13Ta–4?6Zr are sensitive to aging temperature. Aging between 500uC and 350uC results in a phase precipitation, while the isothermal v phase forms during aging between 300uC and 400uC. Ti–39Nb–13Ta–4?6Zr is more stable and no definite evidence of a phase was observed during aging up to 48 h. 3. The cooling rate following solution treatment of Ti–29Nb–13Ta–4?6Zr has a profound influence on phase transformation in the temperature range when a and v coprecipitate from the b matrix. The a0 martensite formed during ice water quenching acted as precursor for the equilibrium a phase and thus provided a low energy path for a precipitation, at the expense of v formation. The resulting dependence of phase fraction and microstructrue on cooling history produces different mechanical properties. 4. Young’s modulus is sensitive to phase fractions in Ti–29Nb–13Ta–4?6Zr and varies significantly with aging temperature. The properties of Ti–39Nb–13Ta–4?6Zr are not responsive to the aging treatment due to more sluggish precipitation reactions.

Acknowledgements a Young’s modulus; b ultimate tensile strength; c yield strength; d elongation 9 Variations of mechanical properties with aging temperature of the two Ti–Nb–Ta–Zr alloys solution treated (WQ) and then aged for 2 days

The research was partially sponsored by Chinese MoST grant TG2000067105 and NSFC grant 50128101. The work of SJL during his visit to Toyohashi was supported by a studentship from the Japanese Ministry of Education and Culture. MN wishes to acknowledge partial support from NEDO and JSPS.

a Young’s modulus; b ultimate tensile strength; c yield strength; d elongation 10 Variations of mechanical properties with cooling rate and aging temperature of Ti–29Nb–13Ta–4?6Zr solution treated (WQ, AC, FC) and then aged for 2 days


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References 1. P. Slanina, W. Frech, A. Bernhardson, A. Cedergren and P. Mattsson: Acta Pharmacol. Toxicol., 1985, 56, 331–336. 2. A. R. Dujovne, J. D. Bobyn, J. J. Krygier, J. E. Miller and C. E. Brooks: J. Arthroplasty, 1993, 8, 7. 3. K. Mishra, J. A. Davidson, R. A. Poggie, P. Kovacs and T. J. Fitzgerald: in ‘Medical application of titanium and its alloys: the material and biological issues’, (ed. J. E. Lemons), ASTM STP 1272, p. 96; 1996, West Conshohocken, PA, ASTM. 4. M. Long and H. J. Rack: Biomaterials, 1998, 19, 1621–1639. 5. M. Niinomi: Mater. Sci. Eng. A, 1998, A243, 231–236. 6. D. L. Moffat and D. C. Larbalestier: Metall. Trans. A, 1988, 19A, 1677–1686. 7. D. L. Moffat and D. C. Larbalestier: Metall. Trans. A, 1988, 19A, 1687–1694. 8. D. L. Moffat and U. R. Kattner: Metall. Trans. A, 1988, 19A, 2389–2397. 9. T. Ahmed and H. J. Rack: J. Mater. Sci., 1996, 31, 4267–4276. 10. X. Tang, T. Ahmed and H. J. Rack: J. Mater. Sci., 2000, 35, 1805–1811. 11. Y. L. Hao, M. Niinomi, D. Kuroda, K. Fukunaga, Y. L. Zhou, R. Yang and A. Suzuki: Metall. Mater. Trans. A, 2002, 33A, 3137–3144. 12. Y. L. Hao, M. Niinomi, D. Kuroda, K. Fukunaga, Y. L. Zhou, R. Yang and A. Suzuki: Metall. Mater. Trans. A, 2003, 34A, 1007–1012. 13. S. J. Li, Y. L. Hao, R. Yang, Y. Y. Cui and M. Niinomi: Mater. Trans., 2002, 43, 2964–2969. 14. ‘Introduction of Young’s modulus measure system (JE–RT 3)’, Technical Manual, p. 7; Japan, Nippon Techno–Plus Co. Ltd.

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15. A. T. Balcerzak and S. L. Sass: Metall. Trans., 1972, 3, 1601–1605. 16. L. Bendersky, W. J. Boettinger, P. B. Burton, F. S. Biancaniello and C. B. Shoemaker: Acta. Metall., 1990, 38, 931–943. 17. S. L. Lopata and E. B. Kula: Trans. Met. Soc. AIME, 1965, 233, 288–293. 18. M. Ikeda, S. Y. Komatsu. I. Sowa and M. Niinomi: Metall. Mater. Trans. A, 2002, 33A, 487–493. 19. T. W. Duerig, R. W. Middleton, G. T. Terlinde and J. C. Williams: in ‘Titanium’80: science and technology’ (eds. H. Kimura and O. Izumi), 1503–1512; 1980, Warrendale, PA, TMS. 20. D. de Fontaine: Acta Metall., 1970, 18, 275–279. 21. S. K. Sikka, Y. K. Vohra and R. Chidambaram: Prog. Mater. Sci., 1982, 27, 245–310. 22. Y. K. Vohra and P. T. Spencer: Phys. Rev. Lett., 2001, 86, 3068–3071. 23. D. Banerjee: Acta Metall. Mater., 1988, 36, 871. 24. R. R. Boyer and G. Lu¨tjering: in ‘Advances in the science and technology of titanium alloy processing’, (ed. I. Weiss et al.), 349–67; 1997, Warrendale, PA, TMS. 25. H. Ledbetter and S. Datta: Z. Metallkd., 1992, 83, 195–198. 26. Z. Fan, P. Tsakiropoulos and A.P. Miodownik: Mater. Sci. Technol., 1993, 9, 863–868. 27. Z. Fan, A. P. Miodownik and P. Tsakiropoulos: Mater. Sci. Technol., 1993, 9, 1094–1100. 28. R. Boccaccini: Z. Metallkd., 1997, 88, 23–26. 29. Y. T. Lee, M. Peters and G. Welsch: Metall. Trans. A, 1991, 22A, 709–714. 30. G. Welsch and W. Bunk: Metall. Trans. A, 1982, 13A, 889–899. 31. Z. Fan: Scr. Metall. Mater., 1993, 29, 1427–1432. 32. E. W. Collings: ‘The physical metallurgy of titanium alloys’, 1984, Metals Park, OH, ASM.