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Materials Science and Engineering A 438–440 (2006) 839–843

Effect of thermo-mechanical treatment on mechanical properties andshape memory behavior of Ti–(26–28) at.% Nb alloys

H.Y. Kim a, J.I. Kim a, T. Inamura b, H. Hosoda b, S. Miyazaki a,∗a Institute of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, Japan

b Precision and Intelligence Laboratory, Tokyo Institute of Technology, Yokohama 226-8503, Japan

Received 13 June 2005; received in revised form 10 January 2006; accepted 6 February 2006

bstract

The effect of thermo-mechanical treatment on mechanical properties and shape memory behavior of Ti–(26–28) at.% Nb alloys was investigatedn order to develop biomedical shape memory alloys. For the solution treated specimens, superelastic behavior was observed in the temperatureanges between 293 and 313 K in Ti–26 at.% Nb, 193 and 313 K in Ti–27 at.% Nb and 163 and 233 K in Ti–28 at.% Nb alloys. However, perfectuperelastic behavior with a strain larger than 2% could not be obtained at room temperature because of the low critical stress for slip deformationn the solution treated alloys. Low temperature annealing (at 873 K) increased the critical stress for slip and stabilized the superelastic behavior.n aging treatment at 573 K after annealing at 873 K further improved the superelastic properties of the Ti–(26–28) at.% Nb alloys. Both the

ensile strength and the critical stress for inducing the martensitic transformation increased with increasing aging time. However, the elongation

ecreased with increasing aging time. The increase of the tensile strength and the aging embrittlement were due to the formation of thermal �hase. The aging effect increased with decreasing Nb content. In particular, perfect superelastic behavior was obtained with the strain up to 3% inhe Ti–26 at.% Nb alloy annealed at 873 K followed by aging at 573 K for 3.6 ks.

2006 Elsevier B.V. All rights reserved.

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eywords: Ti-base alloys; Biomedical shape memory alloys; Superelasticity; T

. Introduction

Ti–Ni shape memory alloys have been used as biomedicalaterials owing to their excellent mechanical properties, high

orrosion resistance and superior superelasticity. Although thei–Ni alloys have been successfully applied for many medicalroducts, the development of Ni-free shape memory alloys istrongly required because of Ni-hypersensitivity and toxicity ofi. Recently the �-type Ti alloys have attracted attention as newiomedical shape memory and superelastic materials [1–12].he �-type Ti alloys exhibit a martensitic transformation from(disordered BCC) to hexagonal martensite (�′) or orthorhom-ic martensite(�′′) depending on alloy composition. The shapeemory effect and superelastic behavior have been reported in

i–Nb, Ti–Mo and Ti–V based alloys. It has been confirmed that

he reversion of �′′ to � results in shape memory behavior in the-type Ti alloys. The present authors have paid attention to Mo

∗ Corresponding author. Tel.: +81 29 853 5283; fax: +81 29 853 5283.E-mail address: miyazaki@ims.tsukuba.ac.jp (S. Miyazaki).

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921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.02.136

o-mechanical treatment; Ti–Nb

nd Nb as a �-stabilizer [6–14] because of cytotoxicity of V.e have reported that the Ti–(22–27) at.% Nb alloys exhibited

hape memory effect and superelastic behavior at room temper-ture [11]. But, the low critical stress for slip deformation causedhe superelasticity not to reveal a large strain in binary Ti–Nblloys.

In this study, effect of thermo-mechanical treatmentn mechanical properties and shape memory behavior ofi–(26–28) at.% Nb alloys was investigated in order to developiomedical shape memory alloys. The shape memory propertynd microstructure were investigated by cyclic tensile tests andtransmission electron microscope.

. Experimental procedure

The Ti–(26–28) at.% Nb alloys were prepared by the Ar arcelting method. The ingots were homogenized at 1273 K for

.2 ks and cold-rolled with the reduction of 95 or 99% in thick-ess. Specimens for the mechanical tests were cut from the cold-olled sheet by an electro-discharge machine. The specimensere cleaned with ethanol, wrapped in Ti foils and encapsulated

8 d Engineering A 438–440 (2006) 839–843

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40 H.Y. Kim et al. / Materials Science an

n quartz tubes under a 3.3 kPa partial pressure of high-purity Ar,nd then heat treated at 1173 K for 3.6 ks or at 873 K for 0.6 ks.ome specimens were aged at 573 K for various times. The spec-

mens were quenched into water by breaking the quartz tubes.he oxidized surface was removed by mechanical polishing fol-

owed by electro-polishing. Tensile tests were carried out at atrain rate of 1.67 × 10−4 s−1 at various temperatures. The gageength of the specimens was 20 mm. Specimens for transmissionlectron microscopy (TEM) observation were also prepared byconventional twin-jet polishing technique. TEM studies were

onducted using a JEOL2010F microscope operated at 200 kV.

. Results and discussion

Fig. 1 shows a series of stress–strain curves obtained at var-ous temperatures for the Ti–(26–28) at.% Nb alloys after theolution treatment at 1173 K for 3.6 ks. The specimens whicho not exhibit complete superelastic recovery upon unloadingere heated up to about 500 K: broken lines with an arrow indi-

ate the shape recovery by heating. The shape memory effect wasbserved for the Ti–26 at.% Nb alloy deformed at temperaturesetween 193 and 273 K. Complete shape recovery occurred byeating to about 500 K. Superelastic behavior was observed at93 and 313 K although the shape recovery was incomplete. Theuperelastic strain decreased with further increasing test temper-ture. The critical stress for apparent yielding corresponds to theritical stress to induce the martensite in the temperature rangef 273–333 K since the martensitic transformation start temper-ture (Ms) of the Ti–26 at.% Nb is about 273 K [11]. For thei–27 at.% Nb alloy, incomplete superelastic behavior was alsobserved at temperatures between 193 and 313 K. The residual

train was recovered by heating. This indicates that the finishemperature of the reverse transformation (Af) is higher than13 K in the Ti–27 at.% Nb alloy. The Ti–28 at.% Nb alloy exhib-ted excellent superelastic behavior at temperatures between 163

ig. 1. Stress–strain curves obtained upon loading and unloading at variousemperatures for the Ti–(26–28) at.% Nb alloys.

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ig. 2. Effect of annealing temperature on stress–strain curve of the Ti–26 at.%b alloy.

nd 233 K. The residual strain increased with increasing tem-erature. Shape recovery was hardly observed when the testemperature was higher than 273 K. The critical stress to inducehe martensite increased with increasing temperature, while theritical stress for slip decreased with increasing temperature.lip occurs if the critical stress level for slip becomes lower

han the stress to induce the martensite. Thus, the strain by slipeformation increased with increasing temperature, causing theecoverable strain to decrease. As a result, it is concluded thathe low critical stress for slip deformation caused the superelas-icity not to reveal a large strain in the solution treated binaryi–Nb alloys.

Fig. 2 shows the stress–strain curves obtained at room tem-erature for the Ti–26 at.% Nb alloy after annealing in theemperature range of 573–1173 K for 3.6 ks. Premature fail-re was observed in the specimens annealed in the temperatureange of 573–673 K. The specimen annealed at 773 K fracturedfter yielding. The specimens annealed in the temperature rangef 873–1173 K exhibited a two-stage yielding. The ultimateensile strength (UTS) decreased with increasing annealing tem-erature. The fracture strain steeply increased with increasingnnealing temperature from 773 to 1173 K. The Ti–26 at.% Nblloy becomes � single phase when the alloy is heat treated at73 K and above, since the � transus temperature is about 823 K.owever, it is noted that the yield stress for plastic deforma-

ion, i.e. the second stage yield stress, decreased with increasingnnealing temperature from 873 to 1173 K. Furthermore, theracture strain increased with increasing heat treatment temper-ture from 873 to 973 K. The differences in the stress–strainurves for the Ti–26 at.% Nb alloy are mainly due to the changen the microstructure such as dislocation density and grain size.

bimodal grain structure consisting of fine subgrains and largeecrystallized grains was observed in the specimen heat treatedt 873 K for 3.6 ks. On the other hand, a fully recrystallizedrain structure was observed in the specimens heat treated at73 and 1173 K. The decrease in the yield stress with increas-ng heat treatment temperature from 973 and 1173 K is due tohe increase in grain size. This indicates that the low tempera-

ure annealing is effective for increasing the critical stress forermanent deformation. Also it has been reported that aging at73 K is effective to increase the critical stress for the solutionreated Ti–(25–27) at.% Nb alloys [11].

H.Y. Kim et al. / Materials Science and Engineering A 438–440 (2006) 839–843 841

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ig. 3. Stress–strain curves obtained by cyclic loading–unloading tensile testst 573 K for (b) 1.8 ks, (c) 3.6 ks and (d) 36 ks after annealing.

In order to investigate the shape memory behavior, strainncrement cyclic tensile tests were carried out for the Ti–26 at.%b specimen subjected to annealing at 873 K for 0.6 ks, and the

esult is shown in Fig. 3. Fig. 3 also shows stress–strain curves ofhe specimens aged at 573 K after annealing at 873 K for 0.6 ks.t the first cycle, tensile stress was applied until strain reached

bout 1.5%, and then the stress was removed in the specimen.he similar measurement was repeated by increasing the max-

mum strain by 0.5% upon loading using the same specimen.he perfect superelasticity was obtained at the first cycle in thepecimen annealed at 873 K. Complete shape recovery occurredy heating at the second and third cycles. As increasing ten-ile strain, the superelastic behavior became incomplete and theemained plastic strain increased.

The specimen aged at 573 K for 1.8 ks exhibited almost per-ect superelasticity at the first and second cycles. The remainedlastic strain increased with increasing applied strain althoughhe superelastically recovered strain increased. The specimenged for 3.6 ks revealed excellent superelasticity. Perfect supere-asticity was observed until the fourth cycle. The maximumuperelastic strain of 3.3% was obtained at the fifth cycle. Its obvious that the critical stress for the first yielding, whichorresponds to the stress for inducing the martensitic transfor-ation, increased when compared with the other two speci-ens. The maximum stress reached at each cycle also increasedith increasing aging time. It is also noted that the stable

uperelasticity was observed even though the maximum stresseached 660 MPa upon loading. On the other hand, the spec-men aged for 36 ks fractured at the third cycle. It has been

ell known that the � phase is Ti-rich although the exact com-osition of � phase is difficult to be determined. This implieshat the precipitation of thermal � phase increases the Nb con-ent of matrix, resulting in decrease of the martensitic trans-

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) the Ti–26 at.% Nb alloy annealed at 873 K for 0.6 ks and the specimens aged

ormation temperature. Furthermore, the dispersed � particlesechanically suppress the martensitic transformation, also caus-

ng the transformation temperature to decrease. As a result,t is reasonable that the stress for inducing the martensiticransformation increased by the precipitation of the thermal

phase. It is also noted that the second stage yield stressncreases with increasing aging time, indicating that the thermal

phase is also effective for increasing the critical stress for slipeformation.

Fig. 4(a) represents a dark field TEM micrograph and theorresponding selected area diffraction pattern of the specimenged at 873 K for 0.6 ks. The selected area diffraction patternas obtained from the [1 1 0]β zone axis. In addition to the pri-ary reflections from � matrix, diffuse scattering at 1/3 {1 1 2}

ositions corresponding to the athermal � phase is visible inhe selected area diffraction pattern. Very fine � particles with

dimension of 3 nm were observed in the dark field TEMicrograph. These � particles are considered as the athermalphase which was formed during quenching after annealing.

he selected area diffraction patterns and dark field TEM micro-raphs obtained from the aged specimens were also shown inig. 4. As increasing aging time, the reflections from � phaseecame discrete spots and the size and volume fraction of �articles increased as shown in the dark field images. The � par-icles larger than 20 nm were observed in the specimen aged at73 K for 36 ks although the size distribution of the specimenas quite broad. This indicates that the substantial increase of

ensile strength and improvement of superelastic property afterging at 573 K for 3.6 ks are due to the growth of thermal �

hase. It is also noted that aging at 573 K for 36 ks resulted in aevere loss of ductility. In conclusion, it is suggested that the finend dense � precipitates are effective for improving superelasticroperty of the binary Ti–Nb alloys.

842 H.Y. Kim et al. / Materials Science and Engineering A 438–440 (2006) 839–843

Fig. 4. Dark-field TEM micrographs and the corresponding selected area diffractionsubsequently aged at 573 K for (b) 1.8 ks, (c) 3.6 ks and (d) 36 ks.

Fig. 5. Stress–strain curves obtained by cyclic loading–unloading tensile testsfor the Ti–27 at.% Nb and Ti–28 at.% Nb alloys annealed at 873 K for 0.6 ks andthe specimens subsequently aged at 573 K for 3.6 ks.

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patterns of (a) the specimen annealed at 873 K for 0.6 ks and the specimens

Similar cyclic tensile tests were carried out for the Ti–27 at.%b and Ti–28 at.% Nb alloys subjected to annealing at 873 K

or 0.6 ks and the result is shown in Fig. 5. Fig. 5 also showstress–strain curves of specimens annealed at 873 K for 0.6 ksollowed by aging treatment at 573 K for 3.6 ks. It can be seenhat almost perfect superelasticity was observed until the sec-nd cycle in the Ti–27 at.% Nb and Ti–28 at.% Nb specimensnnealed at 873 K for 0.6 ks. The apparent yield stress increasedith increasing Nb content. This is due to that the Ms temperatureecreases with increasing Nb content, causing the critical stressor inducing the martensitic transformation to increase. Moretable superelasticity was obtained after aging at 573 K for 3.6 ksn both alloys. It is also noted that the increment of the apparentield stress by aging decreased with increasing Nb content. Thiss reasonable because the stability of � increases with increasingb content [15]. Based on the above results, it is concluded that

uperelastic property of the Ti–Nb alloys can be improved byhermo-mechanical treatment and the Ti–(26–28) at.% Nb alloysre promising for the biomedical superelastic alloys.

. Conclusions

(i) The shape memory effect and/or superelastic behavior wereobserved in the solution treated Ti–(26–28) at.% Nb alloys.The superelastic behavior was observed in the tempera-

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H.Y. Kim et al. / Materials Science an

ture ranges between 293 and 313 K in the Ti–26 at.% Nb,193 and 313 K in the Ti–27 at.% Nb and 163 and 233 Kin the Ti–28 at.% Nb alloys, respectively. However, perfectsuperelastic behavior with a strain larger than 2% could notbe obtained at room temperature because of the low criticalstress for slip deformation in the solution treated alloys.

(ii) Ultimate tensile strength decreased and fracture strainincreased with increasing annealing temperature. The lowtemperature annealing at 873 K stabilized the superelastic-ity in the Ti–(26–28) at.% Nb alloys.

iii) The aging treatment at 573 K after annealing at 873 Kincreased both of the tensile strength and the critical stressfor inducing the martensitic transformation due to the for-mation of thermal � phase in the Ti–(26–28) at.% Nballoys. However, the elongation decreased with increasingaging time. Perfect superelastic behavior was obtained withthe strain up to 3% by annealing at 873 K followed by sub-sequent aging at 573 K for 3.6 ks in the Ti–26 at.% Nb alloy.

cknowledgments

This work was partially supported by ILC Project from Uni-

ersity of Tsukuba and the 21 Century Center of Excellence Pro-ram and the Grants-in-Aid for Fundamental Scientific ResearchKiban A (1999–2001), Kiban A (2002–2004) from the Ministryf Education, Culture, Sports, Science and Technology, Japan.

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