Materials Science and Engineering C 33 (2013) 1629–1635

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Materials Science and Engineering C

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Mechanical properties and microstructures of β Ti–25Nb–11Sn ternary alloy forbiomedical applications

Taek-Kyun Jung 1, Satoshi Semboshi, Naoya Masahashi, Shuji Hanada ⁎Institute for Materials Research, Tohoku University, Sendai 980–8577, Japan

⁎ Corresponding author at: Institute for Materials ResKatahira, Aoba-ku, Sendai 980–8577, Japan. Tel./fax: +

E-mail address: hanada@imr.tohoku.ac.jp (S. Han1 Present address: Korea Institute of Industrial Techn

lic of Korea.

0928-4931/$ – see front matter © 2013 Elsevier B.V. Allhttp://dx.doi.org/10.1016/j.msec.2012.12.072

a b s t r a c t

a r t i c l e i n f o

Article history:Received 7 June 2012Received in revised form 14 November 2012Accepted 18 December 2012Available online 23 December 2012

Keywords:β titanium alloyYoung's modulusStrengthα precipitation

The mechanical properties and microstructures of β Ti–25%Nb–11%Sn ternary alloy rods were investigatedfor biomedical applications as a function of heat treatment temperature after swaging by an 86% reductionin cross-section area. An as-swaged rod consisting of a β (bcc) single phase shows a low Young's modulusof 53 GPa, which is interpreted in terms of both the metastable composition of the β alloy undergoing neitheran athermal ω transformation nor a deformation-induced ω transformation and b110>texture developmentduring swaging. Heat treatment at 673 K (400 °C) for 2 h leads to a high strength of approximately 1330 MPaand a high spring-back ratio of yield stress to Young's modulus over 15×10−3, with acceptable elongation.This high strength is attributable to needle-like α precipitates, which are identified by high-angle annulardark-field scanning transmission electron microscopy (HAADF-STEM) and high-resolution electron micros-copy (HREM).

© 2013 Elsevier B.V. All rights reserved.

1. Introduction

Currently used orthopedic implant alloys such as Ti–6Al–4V andCo–Cr have Young's moduli greater than 100 GPa [1], while bonemoduli are 4–30 GPa, depending on the type of bone and the directionof measurement [2,3]. This large difference between implant devicesand bone tissues causes bone resorption due to stress shielding [4], andit results in the loosening of devices,making total hip arthroplasty neces-sary. To overcome the critical issue of stress shielding, β Ti–Nb-based al-loys for biomedical applications have recently attracted much attention,andmany alloys with low Young's modulus as well as high strength andgood biocompatibility have been developed, e.g., Ti–13Nb–13Zr [5],Ti–35Nb–5Ta–7Zr [6], Ti–29Nb–13Ta–4.6Zr [7] and Ti–33.6Nb–4Sn[8], etc. These alloys possess low Young's moduli of 40–80 GPa; thus,they have the potential to replace conventional implant alloys. Amongthem, β Ti–Nb–Sn alloys have low melting temperature due to Snalloying and no additional refractory element with a high melting tem-perature, such as Ta. This implies that large ingots can be easily madeand the chemical homogenization of constituent elements is readilyattained duringmelting and subsequent hot working. β Ti–Nb–Sn al-loys have been found to provide attractive mechanical properties, in-cluding low Young's modulus [8–10], high strength [9,10] and good

earch, Tohoku University, 2-1-181 22 215 2117.ada).ology, Incheon 406–840, Repub-

rights reserved.

biocompatibility [11]. Moreover, β Ti–Nb–Sn alloys are very ductileat ambient temperature, which enables cold die-forging to fabricatenear-net-shaped implants [12]. These properties are achieved by opti-mizing the associated chemical compositions and thermo-mechanicalprocesses. It is noted that Sn acts as a β stabilizing element in the βphase of Ti alloys in a manner similar to that of Nb [8,13]; hence,metastable β Ti–Nb–Sn alloys can be designed to possess lowerYoung's moduli over a wide composition range by balancing the Nband Sn contents. For example, metastable Ti–33.6Nb–4Sn produces aslight α″martensite phase in a βmatrix after quenching and noticeablyafter subsequent cold rolling [9]. Previously, Ti–22.3Nb–11Snwas shownto behave similarly as far as constituent phases and their volume frac-tions before and after cold rolling are concerned. Upon further increasingthe Sn content to 12%, α″ is not detected, even after cold rolling inTi–22Nb–12Sn. Matsumoto et al. [10] showed that the mechanicalproperties of thermo-mechanically processed Ti–33.6Nb–4Sn aregreatly influenced by deformation-induced α″ with strong rollingtexture. However, they did not clarify the effect of the texture of thebcc matrix phase on mechanical properties.

The objective of this paper was to investigate the mechanical prop-erties of a metastable β Ti–Nb–Sn alloy undergoing no deformation-induced α″ transformation for comparison with the properties of aless stable β Ti–Nb–Sn alloy. Ti–25Nb–11Sn alloy with a lowNb contentwas selected tomeet this objective because the expensive Nb content isreduced to retain themetastableβ single phase for future commercial ap-plications. Yield stress, ultimate tensile strength, elongation and Young'smodulus were examined in tensile tests as a function of heat treatmenttemperature after cold swaging for the development of β Ti alloy im-plants with low Young's modulus and high strength. Microstructures

Fig. 1. XRD patterns of as-swaged β Ti–25Nb–11Sn rod along sections perpendicularand parallel to the longitudinal direction.

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Fig. 2. Temperature dependence of yield stress, tensile strength and Young's modulusof Ti–25Nb–11Sn heat treated after swaging.

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were investigated mainly by transmission electron microscopy (TEM) tounderstand the effects of heat treatment on themechanical properties ofthe materials studied.

Once a low Young's modulus with a high yield stress is obtained,an alloy can be applied not only in orthopedic implants but also ortho-dontic devices because a large elastic recovery strain arising fromHook's law is expected. This paper also discusses the effect of heat treat-ments on the recovery of elastic strain by evaluating the spring-backratio.

2. Experimental procedure

A Ti–Nb–Sn alloy ingot weighing 15 kg was prepared by high-frequency induction melting in a vacuum using a water-cooled cop-per hearth, followed by vacuum arc remelting. The nominal composi-tion of the alloy was Ti–25%Nb–11%Sn (compositions in this paper aredenoted bymass% in the style of Ti–25Nb–11Sn unless otherwise noted).The ingot was hot forged at 1373 to 1173 K (1100 to 900 °C) to a30 mmϕ rod, hot rolled at 1173 K (900 °C) to an 8 mmϕ rod and swagedat room temperature to a 3 mmϕ rod after removing surface oxide layersby scalping. The swaging reduction was 86% in cross section. The chem-ical composition after swaging was Ti–25.4%Nb–11.2%Sn, with impu-rities of 0.07%O, 0.005%C, 0.007%N and 0.02%Fe. The alloy will bereferred to as Ti–25Nb–11Sn hereafter. Tensile samples with the gageportion of 15 mm in length and 2 mm in diameter were prepared. Ten-sile tests at an initial strain rate of 5.6×10−4 s−1 were carried out at297 K (25 °C) after heat treatments at 473 to 1173 K (200 to 900 °C)for 2 h. At least 3 specimens were heat treated at each temperatureand tensile tested. Young's modulus wasmeasured by recording tensilestrain using a video extensometer.

Constituent phases after swagingwere determined by X-ray diffrac-tion (XRD)with Cu–Kα radiation. Themicrostructures of the alloywereinvestigated using an optical microscope (OM), a scanning electronmicroscope (SEM) and a transmission electron microscope (TEM). βtransus temperature was determined by the differential scanningcalorimetry coupled with microstructural observations. Using a JEOLJEM-3010with a LaB6 gun operating at 300 kV, electron diffraction pat-terns were taken from a selected area of approximately 80 nm in diam-eter. High-angle annular dark-field scanning transmission electronmicroscopy (HAADF-STEM) images were acquired using an FEI Titan80–300 TEM operating at 300 kVwith a field emission gun and a detec-tor with an inner angle greater than 55 mrad. To prepare TEM samples,disks perpendicular to the longitudinal direction of a swaged bar wereground to less than 80 μm in thickness and electro-polished in a solu-tion of 10 vol.% nitric acid in methanol at 243 K (−30 °C) with a DCvoltage below 10 V. Finally, thin foils were prepared using a low-angleion milling machine, Fischione Instruments Model 1010, with anargon ion beam accelerated at 3 kV.

3. Results and discussion

Swaging from 8 to 3 mmϕ was easily carried out at room tempera-ture without intermediate annealing, indicating the excellent ductilityof the alloy. Fig. 1 shows XRD patterns taken from cross sections perpen-dicular and parallel to the longitudinal direction of the swaged rod,whichare denoted as ‘perpendicular’ and ‘parallel’ in the figure, respectively. Itcan be seen that all the reflections are regarded as bcc structure and 110reflection in the perpendicular is very strong, as compared to other re-flections such as 220 and 211. Therefore, Ti–25Nb–11Sn consists of a β(bcc) single phase with a strong b110>β texture along the swaging di-rection. Fig. 2 shows the yield stress (flow stress at 0.2% plastic strain), ul-timate tensile strength and Young's modulus of the as-swaged andheat-treated Ti–25Nb–11Sn as a function of heat treatment temperature.With increasing temperature, all of the mechanical properties increaseand exhibit a peak approximately 673 K (400 °C). The peak yield stressand tensile strength are approximately 1300 MPa and 1330 MPa,

respectively, and decrease with a further increase in temperature.Young'smodulus along the longitudinal direction is 53 GPa after swagingand increases to 86 GPa at the peak temperature. A hot-rolled Ti–25Nb–11Sn rod showedweak b110>β recrystallization texture along the longi-tudinal direction, and its Young's modulus was found to be scattered inthe range of 50 to 65 GPa before swaging, implying that heavy swagingleads to a slight decrease in Young's modulus with reduced datascattering.

The Young's modulus of β Ti–33.6Nb–4Sn was reported to decreaseto less than 45 GPa by cold rolling at ambient temperature. The lowmodulus was explained in terms of a strong anisotropy in the Young'smodulus of deformation-induced orthorhombic α″, with preferred de-formation texture developing during rolling [10]. However, the currentTi–25Nb–11Sn exhibits no detectable deformation-induced α″. Tane etal. [14] demonstrated that a β Ti–Nb–Ta–Zr single crystal shows an an-isotropy in its Young's modulus, and experimentally measured valuesare on the order of [100]βb [110]βb [111]β. Inamura et al. [15] havedrawn the same conclusion from experiments and the calculation ofthe orientation dependence of Young's modulus in Ti–Nb alloy. Quiterecently, Zhang et al. [16] measured Young's moduli of 27.1, 56.3 and88.1 GPa for b100>, b110>and b111>, respectively, using single crys-tals of Ti–24Nb–4Zr–8Sn. These results are consistent with each other,and it is interesting to note that the Young's modulus of b110> in Ti–

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Fig. 3. Temperature dependence of elongation and spring-back ratio of Ti–25Nb–11Snheat treated after swaging.

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24Nb–4Zr–8Sn with the composition similar to Ti–25Nb–11Sn in thisstudy is 56.3 GPa. The observed Young's modulus of 53 GPa after swag-ingwill be associatedwith the strong b110>β texture development ob-served experimentally (Fig. 1). On the other hand, in Ti–33.6Nb–4Sn,the preferred rolling textures of b110>β for β and b020>α″ for α″were observed along the rolling direction, and a Young's moduluslower than 50 GPa was measured [10]. Therefore, the lower Young'smodulus of Ti–33.6Nb–4Sn versus that of Ti–25Nb–11Sn is consideredto be apparent values originating from the anisotropic α″ formation inthe former alloy.

It is widely accepted that a low Young's modulus in β Ti alloys isobtained at the composition containing the least amount of β stabilizingelement to form a β single phase [13,17–19]. This is the case in Ti–25Nb–11Sn, whose composition was experimentally determined by referringto a previous paper [12]. The composition of the prepared alloy is ingood agreementwith the theoretical prediction of phase stability in Ti al-loys made by Morinaga et al. [13,20]. These researchers revealed thatthere is a general correlation between the phase stability and the elasticproperties ofβ Ti alloys by delineating theβ/β+ω phase boundary in theBo–Md diagram. Bo is the bond order, which is ameasure of the covalentbond strength between Ti and an alloying element. Md is the metald-orbital energy level of an alloying transition metal, which correlateswellwith the electronegativity andmetallic radius of elements. The aver-age values ofBo andMd are defined by taking the compositional averageof each parameter. Ti–25Nb–11Sn, which is calculated to have Bo=2.811 and Md=2.424, is located near the Ms=RT curve in the Bo–Mddiagram, where Ms=RT denotes that the alloy has martensitic transfor-mation start temperature around room temperature. Therefore, the lowYoung's modulus of approximately 53 GPa in Ti−25Nb−11Sn can beexplained by the low stability of the β phase.

It should be noted that theBo value of Ti−25Nb alloyedwith Snwitha low Bo value is much smaller than that of Ti−50Ta–20Zr (Bo=2.934)or Ti–30Zr–10Nb–10Ta (mol%, Bo=2.945), although Abdel-Hady et al.predicted that alloying with elements with high Bo values is preferableto reduce Young's modulus [13]. This inconsistency is not clearly under-stood, but the clustermodel of bcc Ti and hcp Ti used in the calculation byMorinaga et al. [20]maybeunapplicable to theTi alloys including so highsolute contents. In addition, some experimental problems may be in-volved in it as follows: 1)β Ti alloys exhibit a crystallographic anisotropywith respect to Young's modulus as described above [14–16], while theyeasily develop recrystallization texture or deformation texture [10].2) Young's modulus in metastable β Ti alloys shows a significant tem-perature dependence at ambient temperature when they possess Msjust below room temperature [15]. 3) The static Young's modulus mea-sured by tensile tests is significantly reduced when the formation of amartensite phase is induced during testing. 4) Young's modulus maybe influenced by the content of unintentionally introduced interstitialimpurity elements such as oxygen and carbon. 5) Young's modulus de-pends on the measurement methods employed; the static Young'smodulus determined by tensile testing is not always equal to the dy-namic one, e.g., via the resonance vibration method. Accordingly, thecomparison of Young's moduli at room temperature using differentpolycrystalline alloys must be performed using the samemeasurementmethod after preparing the samples, through which the above uncer-tainties are avoided as much as possible.

Zener [21]first pointed out that bcc phase stability is associatedwiththe shear modulus c′=(c11–c12)/2 for {110}b110>shear. Fisher andDever [22] showed that c′ decreases with decreasing electron-atomratio (e/a), thereby decreasing the bcc phase stability. Very recently,Lee et al. [23] have revealed that low Young's modulus appears whenβ Ti alloys possess a low value of e/a and ω phase is not included,based on the measurements of the elastic moduli in β Ti–Mo–Zr–Alalloy single crystals and themodulus data of other β Ti alloy single crys-tals in references. The examined β Ti alloys had e/a ratios ranging from4.1 to 4.8. They found that low Young's modulus of 44.4 GPa along theb100> direction is obtained in β Ti–Mo–Zr–Al with the lowest e/a of

4.10.β Ti–25Nb–11Sn alloy in this study has a lowe/a of 4.16. Therefore,the low Young's modulus will result from the low stability of β phase.

Fig. 3 shows the elongation to fracture and spring-back ratio (yieldstress at 0.2% plastic strain to Young's modulus) of as-swaged andheat-treated Ti–25Nb–11Sn. The elongation is 16% after swaging and de-creases to approximately 8% at 623 and 673 K (350 and 400 °C); it in-creases to 30% upon further increasing the temperature. The failure ofthe as-swaged specimen was accompanied with considerable neckingbefore fracture, as shown in Fig. 4(a). The SEM fracture surface corre-sponding to the white line square in Fig. 4(a) is completely composedof dimples (Fig. 4(b)), although the core is different from the rim in con-trast and dimple size. The macroscopic fracture plane is roughly perpen-dicular to the tensile direction at the core,while it is inclined to the tensiledirection at the rim. The SEM fractographs shown in Fig. 4(c and d) for asample fractured at 673 K (400 °C), indicating different fracture modesat the core and the rim, are similar to those shown in Fig. 4 (a and b).These observations suggest that final fracture occurs by void coalescenceunder tensile stress at the core and under shear stress at the rim. Thesefractographic features nearly disappeared when the sample wasfractured at 923 K (650 °C) and completely disappeared at 1073 K(800 °C), as shown in Fig. 4(e, f). Grain growth after recrystallizationwas observed by OM, and no apparent change in average grain sizeswas observed from the rim to core of the rod. Therefore, the characteristicfractographs shown in Fig. 4(a–d) cannot currently be unambiguouslyexplained. Further investigation is necessary. The measured grain sizeafter recrystallization is plotted in Fig. 5, indicating a gradual increasefrom40 μm for 923 K to 117 μmfor 1173 Kwith heat treatment temper-ature aboveβ transus of 925 K (652 °C). Theβ grain growth is consideredto contribute to the decrease in strength observed at high temperatures(Fig. 2). Thus, it was found that the β Ti–25Nb–11Sn alloy is extremelyductile and significantly strengthened by heat treatment at approximate-ly 673 K (400 °C) after swaging.

Fig. 6 shows a TEM bright-field image of the as-swaged sample ob-served along the longitudinal direction.Most of thefine grainsmeasuring100 to 500 nm exhibit remarkable strain contrasts, most likely due to ahigh density of dislocations. All of the grains observed were determinedto possess a bcc structure (lattice parameter, a=0.331 nm) by indexingthe selected area diffraction patterns (an example is inserted in Fig. 6),consistent with the XRD result described above (Fig. 1). It is noted inthe diffraction pattern of Fig. 6 that neither a diffraction spot nor diffusescattering related to theω phase is observed. The formation of athermalω or deformation-inducedω is well known to increase Young'smodulus.Therefore, the low Young's modulus of the as-swaged alloy will be relat-ed tometastableβ Ti alloy containing noω phasewith b110>β deforma-tion texture.

Fig. 4. SEM fractographs for Ti–25Nb–11Sn as-swaged (a,b) and heat treated at 673 K (c,d) and 1073 K (e,f).

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TEM images obtained at the peak strength are shown in Fig. 7,where Fig. 7(a) is a bright-field image and Fig. 7(b) and (c) are typicalselected area diffraction patterns at incident beams of [001]β and[11—0]β. All grains in Fig. 7(a) contain fine needle-like products as

well as strain contrasts, most likely arising from a high density of dis-locations. The diffraction patterns in Figs. 7(b) and (c) show weak re-flections, together with strong fundamental spots from bcc structure.The weak reflections are from α phase with a distorted hcp structure(a=0.297 nm, c=0.471 nm) as indexed in thefigures and streaks seenin Fig. 7(c) are probably due to a shape-effect by thin needle α phase.The arrowed reflection maxima are close to, but slightly displacedfrom the expected positions of ω diffraction maxima. The reflectionswould come from α precipitates with other variants, although theywere not identified unambiguously because of weak reflections. It hasbeen revealed that the ω transformation is retarded or suppressed byadding a small amount of elements such as Sn, Al, Zr and O to metasta-ble β Ti alloys [13,17,24–30]. Therefore, the suppression of thermal ωphase precipitation in the present alloy may be due to the relatively

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Fig. 5. Temperature dependence of β grain size for Ti–25Nb–11Sn heat treated afterswaging.

high Nb and/or Sn content. Fig. 8(a) shows a high-magnification TEMbright-field image at the peak strength. Fine needle-like precipitatesof the α phase, measuring 20–100 nm in length, can be observed alongtwo b111>β directions. Fig. 8(b) shows a HAADF-STEM image, whereprecipitates are manifested as the dark contrast. Z contrast imagesobtained by HAADF-STEM are known to result from the thermal diffusescattering of atoms which is approximately proportional to the squareof the atomic number. It is evident from a comparison of the atomicnumber of Nb with that of Ti and Sn that the Nb content in the precipi-tates is lower than that of the β matrix. Hence, the HAADF-STEM resultis consistent with the above conclusion from the electron diffractionanalysis that the precipitates are composed of the α phase formed bythe decomposition of the metastable β phase.

Fig. 6. TEM bright-field image of Ti–25Nb–11Sn as-swaged, where a typical selectedarea diffraction pattern taken from a grain viewed along [11

—0]β is shown in the inset.

Fig. 7. TEM bright-field image of Ti–25Nb–11Sn heat treated at 673 K (a) and electron diffraction patterns viewed along [001]β (b) and [11—0]β (c). Zone axis of the matrix [001]β

with two variants of [1100]α zones in (b).

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Fig. 9 shows a HREM image taken from the [110]β zone axis in anarea containing precipitates at the peak strength. The image showsthat the precipitates with incoherent α/β interface possess a hexago-nal structure whose [0001]α is parallel to the zone axis and b112

—0>α

is parallel to b111>β of the β matrix. Fast Fourier Transform (FFT)images taken from the selected regions of β and α phases are shownin Fig. 9(b) and (c), respectively. The FFT analysis of the HREM image

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10 nm

Fig. 8. TEM bright-field image (a) and Z contrast image obtained

confirms that the Burger's orientation relationship between β and α,{110}β//{0001}α and b1–11>β//b11–20>α, is satisfied in the presentalloy as in many β Ti alloys with α precipitates. It has been reportedthat metastable β Ti–Nb-based alloys are significantly strengthened bythe decomposition of β to α, ω or β′, depending sensitively on thealloy composition and heat treatment condition [25,28,31–36]. More-over, some alloys have been found to be strengthened by B2 ordering

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Fig. 9. HREM image of Ti–25Nb–11Sn heat treated at 673 K (a) and Fast Fourier Transform of regions consisting of β (b) and α (c).

500 nm

Fig. 10. TEM bright-field image of Ti–25Nb–11Sn heat treated at 923 K.

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[37,38]. According to the diffraction patterns shown in Fig. 7(b,c), nodetectable reflections can be observed at the 1/2 {200} positions aris-ing from B2 ordering. Additionally, the splitting of {110}β reflectionsevidencing the coexistence of two phases β and β′ is not observed inthe FFT of Fig. 9(b). Therefore, both β phase separation and B2 orderingcan be ruled out as strengthening mechanisms for age-hardening in βTi–25Nb–11Sn. The preferential precipitation of α over ω, which is re-lated to the stability of the β phase, is not clearly understood at themo-ment, but two factors seem to be involved. One is the high density ofdislocations acting as α nucleation sites in heavily swaged samplesprior to aging; otherwiseω/β interfaces may be needed asα nucleationsites for fine distribution in β grains. The other is the high alloying con-tent of Sn,which acts as an element suppressing or retardingω transfor-mation, as described above.

As shown in Figs. 7, 8 and 9, a high density of fine α needles leadsto a decrease in spacing between the needles, which effectively hin-ders dislocation movement. Young's modulus of Ti alloys is sensitiveto constituent phases and their volume fraction. α phase has higherYoung's modulus than β phase [39]. Thus, the remarkably increasesof Young's modulus, yield stress and tensile strength in Fig. 2 can beascribed to fine needle-like α precipitation.

Fig. 3 also shows the temperature dependence of elastic strain es-timated by the spring-back ratio. A relatively large spring-back ratio over15×10−3 is obtained by aging at 473 K (200 °C) to 723 K (450 °C).Verstrynge et al. [40] evaluated the mechanical properties, includingthe spring-back ratio, of orthodontic archwires produced by β Ti alloysand compared them to those of stainless steel archwires. They dem-onstrated spring-back ratios of (11–14)×10−3 for β Ti alloys and(9–10)×10−3 for stainless steel and reported the advantage of β Tiarchwires with respect to their high spring-back ratio, low stiffness andgood formability. Compare to their results, the results of this study indi-cate that Ti−25Nb−11Sn exhibits a better balance of yield strength andYoung's modulus. However, for a reasonable comparison with their re-sults, the mechanical properties of Ti−25Nb−11Sn should be investi-gated using thin wires of the same size as those they employed.

Fig. 10 shows a bright-field image after heat treatment at 923 K(650 °C), indicating grain growth by recrystallization and a decreasein dislocation density, compared with Fig. 7. Some grains include weakstrain contrast elongated along b111>β, which may be associated withthe α phase, suggesting that β transus temperature is just above 923 K(650 °C). These microstructural changes in substructure may be closelyrelated to the decrease in strength at high temperatures.

As described above, β Ti−25Nb−11Sn alloy has excellent ductilityand low Young's modulus at ambient temperature, and its strength isremarkably increased by α precipitation during heat treatment after

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cold swaging. Therefore, this alloy is adaptable to cold-working orplastic-forming techniques such as cold rolling, drawing, swaging ordie-forging in the fabrication of biomedical devices. Moreover, the me-chanical properties can be controlled by heat treatment after coldwork-ing over the wide range of 900 to 1300 MPa for tensile strength and 53to 86 GPa for Young's modulus.

4. Summary

β Ti−25Nb−11Sn ternary alloy was developed for easy makingof a bulk ingot and easy homogenization of chemical composition,as compared to multicomponent β Ti alloys. The alloy is so ductile atambient temperature that it can be readily swaged to a highly reducedcross section without intermediate annealing. An as-swaged rod has alow Young's modulus of 53 GPa along the longitudinal direction. Thelow Young's modulus is derived from both metastable β Ti alloycontaining neither athermal nor deformation-induced ω phase andpreferred [110]β texture development. With increasing heat treat-ment temperature, the yield stress and tensile strength increaseand exhibit a peak approximately 673 K (400 °C), with a yield stressof 1300 MPa and tensile strength of 1330 MPa. Large elastic strain(spring-back ratio) is obtained by aging around the peak tempera-ture. HAADF-STEM and HREM revealed that strengthening is due tofine needle-like α precipitates.

It is suggested that biomedical devices with attractive mechanicalproperties are easily fabricated by conventional cold working or plas-tic forming processes followed by heat treatment.

Acknowledgments

This workwas supported by the Japan Science and Technology Agen-cy. The authorswould like to thankMessrs. E. Aoyagi and Y. Hayasaka fortheir helpful assistance in TEM observations.

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