ET

Ha

b

c

d

ARRAA

KSTMD

1

msmioTott2mporttot

T

m

0d

Journal of Alloys and Compounds 577S (2013) S435–S438

Contents lists available at ScienceDirect

Journal of Alloys and Compounds

journa l homepage: www.e lsev ier .com/ locate / ja l l com

ffect of Nb content on deformation behavior and shape memory properties ofi–Nb alloys

. Tobea, H.Y. Kima,∗, T. Inamurab, H. Hosodab, T.H. Namc, S. Miyazakia,c,d,∗∗

Division of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, JapanPrecision and Intelligence Laboratory, Tokyo Institute of Technology, Yokohama 226-8503, JapanSchool of Materials Science and Engineering & ERI, Gyeongsang National University, 900 Gazwadong, Jinju, Gyeongnam 660-701, Republic of KoreaCenter of Excellence for Advanced Materials Research, King Abdulaziz University, P.O. Box 80203, Jeddah 21589, Saudi Arabia

a r t i c l e i n f o

rticle history:eceived 31 October 2011eceived in revised form 20 January 2012ccepted 1 February 2012

a b s t r a c t

Deformation behavior and shape memory properties of Ti–(20, 23) at.% Nb alloys in a single �′′ martensitestate were investigated. The Ti–20Nb alloy exhibited a higher stress for the reorientation of martensitevariants when compared with the Ti–23Nb alloy. The recovery strain due to the shape memory effect inthe Ti–20Nb alloy was smaller than that in the Ti–23Nb alloy. Transmission electron microscope (TEM)

vailable online 15 February 2012

eywords:hape memory alloyi–Nbartensite

observation revealed that the reorientation of martensite variants occurred by the deformation of {1 1 1}type I and 〈2 1 1〉 type II twins. The Nb content dependence of the deformation behavior and shape memoryproperties was discussed considering the magnitude of twinning shear of the twins.

© 2012 Elsevier B.V. All rights reserved.

eformation twin

. Introduction

Recently, �-type Ti–Nb based alloys consisting of non-toxic ele-ents have attracted considerable interest as Ni-free biomedical

hape memory alloys [1–10]. These alloys undergo a thermoelasticartensitic transformation from � phase to �′′ phase upon cool-

ng. The crystal structure of the � phase is disordered bcc and thatf the �′′ phase is disordered C-centered orthorhombic [11,12].he martensitic transformation start temperature (Ms) dependsn the alloy composition. For the case of Ti–Nb binary alloys,he Ms increases with decreasing Nb content and becomes higherhan room temperature (RT) when the Nb content is lower than5.5 at.% [1]. Furthermore, it has been reported that the transfor-ation strain calculated from the lattice parameters of the � and �′′

hases increases with decreasing Nb content [1]. The shape mem-ry effect of the alloys in a single �′′ martensite state is due to theeorientation of martensite variants during loading and the reverseransformation upon heating after unloading. Thus the investiga-

ion of the deformation behavior, i.e., the reorientation mechanismf the �′′ martensite variants is important for the understanding ofhe shape memory properties of �-type Ti-based alloys. However,

∗ Corresponding author. Tel.: +81 29 853 6942; fax: +81 29 853 6942.∗∗ Corresponding author at: Division of Materials Science, University of Tsukuba,sukuba, Ibaraki 305-8573, Japan. Tel.: +81 29 853 5283; fax: +81 29 853 5283.

E-mail addresses: heeykim@ims.tsukuba.ac.jp (H.Y. Kim),iyazaki@ims.tsukuba.ac.jp (S. Miyazaki).

925-8388/$ – see front matter © 2012 Elsevier B.V. All rights reserved.oi:10.1016/j.jallcom.2012.02.023

little has been reported on the effect of Nb content on the deforma-tion behavior and shape memory properties of Ti–Nb alloys in thesingle �′′ martensite state. In this study, deformation behavior andshape memory properties of Ti–20 at.% Nb and Ti–23 at.% Nb alloyswere investigated. The Nb content dependence of the deformationbehavior and shape memory properties was discussed on the basisof the results obtained by tensile tests and transmission electronmicroscope (TEM) observation.

2. Experimental

Ti–20 at.% Nb and Ti–23 at.% Nb alloy ingots were fabricated using an Ar arc melt-ing method. The ingots were sealed in a vacuumed quartz tube and homogenizedat 1273 K for 7.2 ks and then quenched into water. The homogenized ingots werecold-rolled at RT with a reduction of 98.5% in thickness. The final thickness of thesheets was about 150 �m. Specimens for X-ray diffraction (XRD) measurements andtensile tests were cut by an electro-discharge machine from the sheets. The oxidizedsurface was removed by chemical etching in a solution of 10 vol.% HF, 40 vol.% HNO3

and 50 vol.% H2O at RT. All the specimens were solution-treated at 1173 K for 1.8 ksin an Ar atmosphere, followed by quenching into water. XRD measurements werecarried out at RT using Cu K� radiation. Tensile tests were conducted at RT along therolling direction at a strain rate of 2.5 × 10−4 s−1. The gage length of the specimenswas 20 mm. Specimens for TEM observation were prepared by a twin-jet polish-ing technique. TEM observation was conducted at RT in a JEOL 2010F instrumentoperated at 200 kV.

3. Results and discussion

Fig. 1 shows the XRD profiles obtained at RT for the Ti–20Nb andTi–23Nb alloys subjected to the solution treatment. In both alloys

S436 H. Tobe et al. / Journal of Alloys and Compounds 577S (2013) S435–S438

Fa

aomrtdaaaa

tTRtltimTsiittmr�c

FT

Table 1Lattice correspondence variants (CVs).

[1 0 0]�′′ [0 1 0]�′′ [0 0 1]�′′

CV1 [1 0 0]� [0 1 1]� [0 1̄ 1]�

CV2 [1 0 0]� [0 1̄ 1]� [0 1̄ 1̄]�

CV3 [0 1 0]� [1 0 1]� [1 0 1̄]�

CV4 [0 1 0]� [1 0 1̄]� [1̄ 0 1̄]�

CV5 [0 0 1] [1 1 0] [1̄ 1 0]

ig. 1. XRD profiles obtained at room temperature for the Ti–20Nb and Ti–23Nblloys.

ll diffraction peaks are indexed as �′′ martensite phase with anrthorhombic structure, indicating that the martensitic transfor-ation finish temperatures of the alloys are higher than RT. This

esult is consistent with previous reports [2,8]. The 2� positions ofhe peaks observed in the Ti–20Nb and Ti–23Nb alloys are slightlyifferent, indicating that the lattice parameters of the �′′ martensitere different from each other. The lattice parameters of the Ti–20Nblloy were determined to be a�′′ = 0.31257 nm, b�′′ = 0.48704 nmnd c�′′ = 0.46456 nm, while those of the Ti–23Nb alloy were

�′′ = 0.31645 nm, b�′′ = 0.48291 nm and c�′′ = 0.46378 nm.Tensile tests were conducted in order to characterize deforma-

ion behavior and shape memory properties of the Ti–20Nb andi–23Nb alloys. Fig. 2 shows the stress-strain curves obtained atT for the solution-treated alloys loaded along the rolling direc-ion. The load was applied until 2.5% strain was reached, then theoad was removed. The specimens were heated after unloadingo measure shape recovery strain induced by heating: the heat-ng temperature is about 600 K, which is higher than the reverse

artensitic transformation finish temperature of the Ti–20Nb andi–23Nb alloys [8]. The dashed lines with an arrow in Fig. 2 repre-ent the strain recovered by heating. Almost perfect shape recoverys observed in the Ti–23Nb alloy, while the Ti–20Nb alloy exhibitsncomplete shape recovery. It is also seen that the yield stress ofhe Ti–20Nb alloy is higher than that of the Ti–23Nb alloy, wherehe yield stress corresponds to the stress for the reorientation of

artensite variants. It should be noted here that the maximum

ecoverable strain, i.e., transformation strain, associated with the–�′′transformation of Ti–Nb alloys increases with decreasing Nbontent [1]. As a result, it is reasonably considered that the smaller

ig. 2. Stress-strain curves obtained at room temperature for the Ti–20Nb andi–23Nb alloys.

� � �

CV6 [0 0 1]� [1̄ 1 0]� [1̄ 1̄ 0]�

recovery strain in the Ti–20Nb alloy is due to the higher yieldstress compared with the Ti–23Nb alloy, because plastic deforma-tion occurs more easily in the alloy with a higher stress for thereorientation of martensite variants upon loading.

The effect of texture should be considered on the deformationbehavior since the crystal orientation affects the stress for the reori-entation of martensite variants and transformation strain. XRD polefigure measurements were conducted at RT for the solution-treatedalloys using diffraction intensities from planes (0 0 2)�′′ , {1 1 1}�′′and (2 0 0)�′′ . There are six lattice correspondence variants (CVs)in the �′′ martensite [13,14] as listed in Table 1. The (0 0 2)�′′ and{1 1 1}�′′ planes correspond to {1 1 0}� planes, whereas the (2 0 0)�′′plane corresponds to {2 0 0}� planes. The texture of the parentphase was determined from the pole figures of the martensitephase based on the lattice correspondence between the martensitephase and the parent phase. The texture formed in the Ti–20Nballoy is {1 1 2}� 〈1 1 0〉�. On the other hand, the texture formedin the Ti–23Nb alloy is composed of a mixture of {1 1 2}� 〈1 1 0〉�and {0 01 }� 〈1 1 0〉�. The transformation strain calculated fromthe lattice parameters of the � and �′′ phases varies according tocrystal orientation [1]. The maximum transformation strain can beobtained when a sample is loaded along 〈1 1 0〉�directions. In boththe Ti–20Nb and Ti–23Nb alloys, one of the 〈1 1 0〉�directions is par-allel to the rolling direction, indicating that the effect of the textureson the transformation strain and stress for reorientation is similar.

In order to investigate the reorientation mechanism of marten-site variants, TEM observation was carried out for the specimensubjected to the solution treatment and for the specimen deformedat RT by a tensile test. Fig. 3(a) and (b) show the bright field imagesof the solution-treated Ti–20Nb alloy, where Fig. 3(b) is the mag-nified image of the framed area in Fig. 3(a). All the six CVs areconfirmed in Fig. 3(a). V-shaped and triangular clusters formed bythick martensite plates were frequently observed. These morpholo-gies have been reported as the self-accommodation morphologiesin the �′′ martensite [13]. The thick martensite plates are mainlyrelated to {1 1 1} type I twinning: the selected area diffraction (SAD)patterns obtained at interfaces between variants CV3 and CV6 andvariants CV2 and CV5 are shown in Fig. 3(c) and (d), respectively.The martensite plates related to 〈2 1 1〉 type II twinning were alsoobserved. Fig. 3(e) and (f) shows the SAD patterns exhibiting the〈2 1 1〉 type II twinning taken from the interfaces between variantsCV6 and CV1 and variants CV6 and CV3, respectively. The twin-ning elements of the {1 1 1} type I and 〈2 1 1〉 type II twins werecalculated by the Bilby–Crocker theory [15] using the lattice param-eters of the �′′ martensite in the Ti–20Nb alloy, and they are listedin Table 2. The K2 and �1 elements of the {1 1 1} type I twin andthe K1 and �2 elements of the 〈2 1 1〉 type II twin have irrationalMiller indices. It is noted that the 〈2 1 1〉 type II twin is the con-jugate of the {1 1 1} type I twin. Fig. 4(a) shows the bright fieldimage of the deformed Ti–20Nb alloy. The Ti–20Nb alloy was loadeduntil the strain leached 2.5% and unloaded, then used for the TEMobservation. The morphology of the deformed martensite is clearly

different from that of the self-accommodated martensite observedin the solution-treated Ti–20Nb alloy (Fig. 3(a)). A nearly singlevariant microstructure formed by the reorientation of martensite

H. Tobe et al. / Journal of Alloys and Compounds 577S (2013) S435–S438 S437

Fig. 3. Bright field images ((a) and (b)) and selected area diffraction patterns with key diagrams exhibiting {1 1 1} type I twinning ((c) and (d)) and 〈2 1 1〉 type II twinning((e) and (f)) of the Ti–20Nb alloy.

Table 2Twinning elements of {1 1 1} type I and 〈2 1 1〉 type II twins calculated from the lattice parameters of the Ti–20Nb alloy.

K1 K2 �1 �2 s⟨ ⟩ ⟨ ⟩

vvfi{aCmtc{sro

twin boundaries between the CVs and/or the introduction of the{1 1 1} type I and 〈2 1 1〉 type II twins in each CV. It is noted that themagnitude of twinning shear of the {1 1 1} type I and 〈2 1 1〉 type II

Table 3Magnitudes of twinning shear of{1 1 1} type I and 〈2 1 1〉 type II twins for the Ti–20Nband Ti–23Nb alloys.

{1 1 1} Type I {1 1 1} {3.201, 5.402, 1}⟨

2 1 1⟩

Type II {3.201, 5.402, 1} {1 1 1}

ariants can be seen in the upper right region of Fig. 4(a). The singleariant grown by the reorientation is identified as CV6. It was con-rmed that the interfaces between the CV6 and the other CVs are1 1 1} type I or 〈2 1 1〉 type II twin boundaries as shown in Fig. 4(b)nd (c). Saburi et al. [16] showed in several �-phase alloys such asu–Al–Ni, Cu–Zn–Ga etc. that the recoverable deformation of shapeemory alloys is produced by mutual conversion among CVs, and

he mechanism of the conversion is twinning deformation. For thease of the �′′ martensite in �-type Ti alloys, deformation of the

1 1 1} type I and 〈2 1 1〉 type II twins is equivalent to the conver-ion of a CV to the other CVs, as reported by Inamura et al. [14]ecently. Therefore, it is considered that the large single variant CV6bserved in Fig. 4(a) seems to be produced by conversion among

5.037, 4.037, 1 2 1 1 0.1613⟨

2 1 1⟩ ⟨

5.037, 4.037, 1⟩

0.1613

the CVs by the movement of the {1 1 1} type I and 〈2 1 1〉 type II

Ti–20Nb Ti–23Nb

{1 1 1} Type I 0.1613 0.1256⟨2 1 1

⟩Type II 0.1613 0.1256

S438 H. Tobe et al. / Journal of Alloys and Compounds 577S (2013) S435–S438

F s tak(

tcstlstphst

4

abtmoTmTcTmita

A

e

[

[

[

[

[Mag. 87 (2007) 3325–3350.

ig. 4. Bright field image (a) and selected area diffraction patterns with key diagramc) of the Ti–20Nb alloy loaded until the strain leached 2.5% and unloaded.

wins depends on lattice parameters of the �′′ martensite, i.e., Nbontent of the Ti–Nb alloys. The calculated magnitudes of twinninghear for the Ti–20Nb and Ti–23Nb alloys are listed in Table 3. Thewinning shears of the two types of twins in the Ti–20Nb alloy arearger than those in the Ti–23Nb alloy. This implies that a highertress and a larger strain are required for deformation by the {1 1 1}ype I and 〈2 1 1〉 type II twinning in the Ti–20Nb alloy when com-ared with the Ti–23Nb alloy. Therefore, it is suggested that theigher stress for the reorientation of martensite variants and themaller recovery strain observed in the Ti–20Nb alloy are due tohe larger shear of the {1 1 1} type I and 〈2 1 1〉 type II twins.

. Conclusions

Deformation behavior and shape memory properties of Ti–20Nbnd Ti–23Nb alloys in a single �′′ martensite state were investigatedy tensile tests and TEM observation. TEM observation revealedhat the reorientation of martensite variants occurred by the defor-

ation of the {1 1 1} type I and 〈2 1 1〉 type II twins. The magnitudef twinning shear in the Ti–20Nb alloy is larger than that in thei–23Nb alloy. The measured recovery strain due to the shapeemory effect in the Ti–20Nb alloy was smaller than that in the

i–23Nb alloy, although a larger transformation strain is expectedonsidering the larger twinning shear strains of the former alloy.he Ti–20Nb alloy exhibited a higher stress for the reorientation ofartensite variants when compared with the Ti–23Nb alloy, result-

ng in larger plastic strain in the former alloy. The higher stress forhe reorientation of martensite variants observed in the Ti–20Nblloy is due to the larger twinning shear.

cknowledgments

This work was partially supported by the Grant-in-Aid for Sci-ntific Research from the Ministry of Education, Culture, Sports,

[[

en form the interfaces between variants CV6 and CV1 (b) and variants CV6 and CV3

Science and Technology, Japan and the Grant-in-Aid for JSPS Fel-lows from the Japan Society for the Promotion of Science. Thiswork was also partially supported by the WCU (World Class Univer-sity) program through the National Research Foundation of Koreafunded by the Ministry of Education, Science and Technology (GrantNumber: R32-2008-000-20093-0).

References

[1] H.Y. Kim, Y. Ikehara, J.I. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 54 (2006)2419–2429.

[2] H.Y. Kim, H. Satoru, J.I. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 45 (2004)2443–2448.

[3] E. Takahashi, T. Sakurai, S. Watanabe, N. Masahashi, S. Hanada, Mater. Trans.43 (2002) 2978–2983.

[4] Y. Fukui, T. Inamura, H. Hosoda, K. Wakashima, S. Miyazaki, Mater. Trans. 45(2004) 1077–1082.

[5] J.I. Kim, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci. Eng. A 403(2005) 334–339.

[6] H.Y. Kim, S. Hashimoto, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Sci.Eng. A 417 (2006) 120–128.

[7] H.Y. Kim, N. Oshika, J.I. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Mater. Trans.48 (2007) 400–406.

[8] Y. Al-Zain, H.Y. Kim, H. Hosoda, T.H. Nam, S. Miyazaki, Acta Mater. 58 (2010)4212–4223.

[9] J.I. Kim, H.Y. Kim, H. Hosoda, S. Miyazaki, Mater. Trans. 46 (2005)852–857.

10] M. Tahara, H.Y. Kim, H. Hosoda, S. Miyazaki, Funct. Mater. Lett. 2 (2009)79–82.

11] A.R.G. Brown, D. Clark, J. Eastabrook, K.S. Jepson, Nature 201 (1964)914–915.

12] T. Inamura, H. Hosoda, H.Y. Kim, S. Miyazaki, Philos. Mag. 90 (2010)3475–3498.

13] Y.W. Chai, H.Y. Kim, H. Hosoda, S. Miyazaki, Acta Mater. 57 (2009)4054–4064.

14] T. Inamura, J.I. Kim, H.Y. Kim, H. Hosoda, K. Wakashima, S. Miyazaki, Philos.

15] B.A. Bilby, A.G. Crocker, Proc. R. Soc. London Ser. A 288 (1965) 240–255.16] T. Saburi, S. Nenno, in: H.I. Aaronson, D.E. Laughlin, R.F. Sekerka, C.M. Wayman

(Eds.), Proc. Int. Conf. on Solid–Solid Phase Transformations, Warrendale, PA,Metall. Soc. AIME (1982) 1455–1479.