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Materials Science and Engineering A 527 (2010) 652–656

Contents lists available at ScienceDirect

Materials Science and Engineering A

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

icrostructure and shape memory effect of Ti–20Zr–10Nb alloy

an Cuia,b, Yan Lia,b,∗, Kun Luoa,b, Huibin Xua,b

School of Materials Science and Engineering, Beihang University, Beijing 100191, ChinaBeijing Key Laboratory for Advanced Functional Materials and Thin Film Technology, Beihang University, Beijing 100191, China

r t i c l e i n f o

rticle history:eceived 20 March 2009eceived in revised form 15 August 2009ccepted 28 August 2009

a b s t r a c t

The microstructure, martensitic transformation behavior and shape memory effect of Ti–20Zr–10Nbshape memory alloy have been studied by X-ray diffraction (XRD), optical microscopy (OM) andtransmission electron microscopy (TEM) observation, differential scanning calorimetry (DSC) and ten-sile stress–strain measurements. The results show that the recrystallization occurs in the cold rolledTi–20Zr–10Nb alloy by annealing at 600 ◦C and the grain size goes up with the increasing annealing tem-

eywords:hape memoryicrostructure

hase transitions

perature up to 800 ◦C. The Ti–20Zr–10Nb alloy is primarily composed of �′′ martensite and a small amountof � phase which appears after annealing. A reverse martensite transformation temperature higher than500 ◦C upon heating has been detected for the Ti–20Zr–10Nb alloy annealed at 600 ◦C, but no obviousexothermic behavior can be found upon cooling. The tensile strength and the failure strain of the alloy

Pa a

i–Zr–Nb alloyold rolling

are measured to be 542 Mof about 2.5%.

. Introduction

Ni-free shape memory alloys (SMAs) based on nontoxic metallements, e.g. Ti–Nb and Ti–Mo, have attracted much attention asdeal candidates for biomedical applications [1–3]. Ti–Nb basedlloys have been extensively investigated on microstructures,hase transformation behaviors, mechanical properties and shapeemory effect (SME) [4–6]. It is known that the shape memory

ffect of Ti–Nb alloys is attributed to the reversible marten-itic transformation between �′′ martensite and parent phase7]. The martensitic transformation temperatures of Ti–Nb alloysecreases with increasing Nb content from the melting point ofi to −90 ◦C [8]. A shape memory effect of 3% has been obtainedn Ti–(22–29 at.%) Nb alloys and the superelasticity strain can bemproved up to 3.3% by cyclic loading–unloading training [2].

Furthermore, some alloying elements such as Sn, Al, Ta, Pd,r have been selected to adjust the martensitic temperature andmprove mechanical and shape memory behaviors of Ti–Nb alloys9–16]. Martensitic transformation temperature of Ti–Nb–Sn alloysecreases rapidly with increasing Sn content and large superelas-

icity strain was obtained in Ti–16Nb–4.9Sn alloy [9], and a shape

emory effect of 4% was reported in a Ti–Nb–Sn alloy [10]. Itas found that both Ta and Al are effective in stabilizing the �hase of Ti–Nb. Moreover, a shape memory effect of 3% and a

∗ Corresponding author at: School of Materials Science and Engineering, Beihangniversity, No.37, Xueyuan Road, Haidian District, Beijing 100191, China.el.: +86 10 82315989; fax: +86 10 82338200.

E-mail address: liyan@buaa.edu.cn (Y. Li).

921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2009.08.063

nd 13.1%, respectively, associated with a maximum shape recovery strain

© 2009 Elsevier B.V. All rights reserved.

large superelasticity of 4.7% have been obtained in Ti–22Nb–4Taalloy [11] and a severe cold rolled Ti–22Nb–3Al alloy [12], respec-tively. Shape memory behavior of a Ti–30Nb–3Pd (wt.%) alloy witha high martensitic transformation temperature of about 173 ◦C wasreported [13].

Recently, Abdel-Hady et al. [14] reported that Zr acts as a �-stabilizer in Ti–Nb–Zr alloys over the wide range of Zr content(6–30 at.%). However, only a few studies have focused on themechanical properties and shape memory behavior of Ti–Nb–Zralloys and these works are related to very low Zr content, e.g.2–6 at.% [15,16]. The composition of Ti–20Zr–10Nb was chosento have � + � two-phase structure (at 500 ◦C) and � singlephasestructure (700 ◦C), respectively, according to the ternary phase dia-grams proposed by Collings [17], which are shown in Fig. 1. In Refs.[14–16], the properties of TiZrNb with low Zr content (the symbol �

as shown in Fig. 1) have been reported, and the Ref. [14] has pointedthat the Nb content necessary for maintaining a � single phase willchange with the addition of Zr to the alloy. The microstructure andshape memory effect of a Ti–Nb based SMA with high Zr contenthave not been focused. So we determine a relative high 20% Zr-containing Ti–20Zr–10Nb alloy (the symbol � as shown in Fig. 1).In the present work, the microstructure and shape memory effectof a Ti–Nb based SMA with high Zr content, Ti–20Zr–10Nb, havebeen investigated, with the effect of cold rolling and annealing inparticular.

2. Experimental

The Ti–20Zr–10Nb (at.%) alloy was prepared from high-puritytitanium, zirconium, and niobium. The ingot was melted in an

Y. Cui et al. / Materials Science and Engineering A 527 (2010) 652–656 653

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Fig. 1. Isothermal section of the Ti–Zr–

rc-melting furnace with a non-consumable W electrode and aater-cooled copper hearth under an ultra-pure argon atmo-

phere. The ingots were turned over and remelted at least fourimes in order to obtain ingots homogeneous in composition and

icrostructure. The button ingots were approximately 1.5 kg ineight. The ingot was cut into 5 mm thick pieces. The pieces of

ast ingot were sealed into evacuated (∼10−3 Pa) quartz tubes andomogenized at 900 ◦C for 2 h followed by quenching into water.nd then deformation by cold rolling under the friction conditionsas carried out at room temperature to a reduction of 80% in thick-ess. The reduction thickness increases from 0.7% to 2.4% per passy successive reduction with an exit velocity of 25 mm s−1. Theolled plates were cleaned by mechanical grinding to remove theurface contaminants, encapsulated in quartz tubes under a vac-um of better than 0.25 Pa, and annealed respectively at 600, 700nd 800 ◦C for 0.5 h, followed by water-quenching. Samples werehen cut from the center section of plates by spark-cutting parallelo the rolling direction to perform the following experiments.

The phase structures present in the microstructure weredentified at room temperature using a Rigaku D/Max 2500C diffractometer operated at 40 kV and 40 mA with a Cu K�adiation (� = 1.5406 nm). Pieces of alloys with average dimen-ions 10 mm × 10 mm × 1 mm for optical microscopy (OM) wereechanically polished and etched at room temperature in a

olution of HF, HNO3 and H2O. Samples for transmission elec-ron microscopy (TEM) were prepared by first grinding to about0 �m and then electro-polishing using the twin-jet method in anlectrolyte of 6% perchloric acid + 35% butanol + 59% methanol at10 ◦C. TEM observations were carried out for microstructure usingJOEL 2010 microscope operated at 200 kV. The phase transfor-ation temperatures of the alloy were determined by differential

canning calorimetry (DSC). A small piece of sample with weightf about 15 mg was cut from the center of rolling plate. In ordero remove the affected surfaces, the samples were mechanicallyolished. DSC experiments were performed on NETZSCH STA409,hich has been calibrated for temperature, enthalpy and base-

ine. Then the samples, protected by Ar gas, were heated up andhen cooled down at a rate of 10 K/min. The tensile stress–strain

xperiments were carried out at room temperature on the MTS80 materials testing system with the tensile direction parallel tohe rolling direction at an initial strain rate of 1.6 × 10−5 s−1, usingensile specimens with a gage section of 1 mm × 1.5 mm × 30 mm.tensile extensometer was used for determining strain during ten-

stem at (a) 500 ◦C and (b) 700 ◦C [17].

sile testing. In order to examine the SME, samples were deformedby tension to a specific strain at room temperature and then heatedto 600 ◦C for strain recovery. The marked lengths of the sampleswere measured before loading (l0), after unloading (l1) and afterheating to 600 ◦C for 3 min (l2) by a micrometer with an accuracyof 0.01 mm. The recovered strain due to the SME were obtained as(l2 − l1)/l0 × 100%. The mechanical behavior data given in this paperare the average of three tensile experiments.

3. Results and discussion

Fig. 2 shows the microstructures of the Ti–20Zr–10Nb alloy. Itcan be seen from Fig. 2a that the as-cast sample exhibits a coarsegrained structure of 300 �m in size, in which lamellar twins arespaced in the order of a few microns. Moreover, the TEM bright fieldimage and selected area electron diffraction (SAED) pattern withthe diffraction spots from [0 0 1]�′′ zone axis for Ti–20Zr–10Nb alloyindicate an internal microstructure of much finer martensitic twinsas identified by Fig. 2b. For the sample with 80% rolling reduction,grain boundaries disappear and martensite plates become obscure,as shown in Fig. 2c. The similar microstructure variation after severecold rolling was also observed in Ti–Nb–Sn alloys [18]. It can be seenfrom Fig. 2d and e that recrystallization occurs for the rolled sampleanealed at 600 ◦C and 700 ◦C. Whereas no recrystallization occursfor the samples annealed at 500 ◦C (not shown). A fine grain with5–10 �m in size is obtained for the sample annealed at 600 ◦C. Thegrain size goes up with increasing annealing temperature, e.g. anaverage value of about 80 �m for the sample annealed at 800 ◦C.In general, grain size has remarkable influence on the mechanicalproperties of alloys; therefore the tensile stress–strain behaviors ofTi–20Zr–10Nb alloy should vary with the annealing temperature asshown later.

Fig. 3 shows the X-ray diffraction patterns of Ti–20Zr–10Nballoys. For the as-cast sample, the metastable martensitic �′′

phase is observed as a predominant phase (C-orthorhombic struc-ture) with the lattice parameters of a = 0.312 nm, b = 0.509 nm andc = 0.479 nm and the b/a ratio as 1.63. This result is similar to thoseof the orthorhombic �′′ martensite in Ti–20Nb–3Al (a = 0.313 nm,

b = 0.485 nm, c = 0.465 nm and b/a = 1.55) [19] and Ti–30Nb–3Pd(a = 0.313 nm, b = 0.482 nm, c = 0.463 nm and b/a = 1.54) [13] alloys.It is seen from curve b in Fig. 3 that after cold rolled, (0 0 2) peakbecomes the most outstanding one and other diffraction peaksbecome weak due to the grain orientation. The similar results have

654 Y. Cui et al. / Materials Science and Engineering A 527 (2010) 652–656

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ig. 2. Microstructures of the Ti–20Zr–10Nb alloy: (a) the optical micrograph and (bs-rolled sample, and as-annealed sample at 600, 700 and 800 ◦C, respectively.

een found in Ti–Nb based alloys [20]. As the annealing tempera-ure rises, the �′′ phase is observed as a predominant phase in eachample annealed at 600, 700, and 800 ◦C. It is noticed that a smallmount diffraction peak of � phase (hexagonal structure) has alsoeen detected in each as-annealed samples. As is well known, thehermo-mechanical process can remodel the microstructure, e.g.igh-pressure distortion can produce � phase in pure Ti [21], andphase also occurs in Ti–Zr alloy after high-temperature and high-

ressure treatment [22]. It has been well-documented that thehermal or aged � phase causes embrittlement [23–25], however,ecent report have shown that the precipitation of fine athermal

phase (about 3 nm in size) does not degrade ductility and isffective in improving the shape memory and superelastic proper-ies in Ti–Nb alloy [2]. Thus, further studies are needed to explorehe effect of � phase on the mechanical properties of Ti–Zr–Nblloy.

TEM bright field image with the inset SEAD pattern for the as-cast sample, (c–f) the

Fig. 4 shows a typical DSC curve of the Ti–20Zr–10Nb alloyannealed at 600 ◦C after cold rolling. Two endothermic peaks onheating can be seen while no transformation peak appears onsubsequent cooling. The first endothermic peak around 300 ◦Cshould correspond to the precipitation of the � phase similar tothat in Ti–Nb–Pd [13] alloys. The second one upon heating cor-responds to the reverse transformation (martensite to austenite),and the transformation temperatures are estimated as the austen-ite transformation peak temperature (Ap) ∼565 ◦C. The martensitictransformation peaks during cooling are too weak to be detected,as indicated in Fig. 4, and similar results have been noticed in

Ti–Nb–Sn [9] and Ti–Nb–Pd [13] alloys. The reason for such a phe-nomenon is unclear now, which might be understood from thefollowing two aspects. Firstly, the enthalpy of the phase transfor-mation between the �′′ martensite and � parent is intrinsic low inTi–20Zr–10Nb alloy. Secondly, a partial transformation from �-Ti

Y. Cui et al. / Materials Science and Engineering A 527 (2010) 652–656 655

Fig. 3. X-ray diffraction pattern of the Ti–20Zr–10Nb alloy at room temperature for:(a) as-cast, (b) as-rolled, (c–e) as-annealed sample at 600, 700 and 800 ◦C, respec-tively.

Fig. 4. DSC curves of the Ti–20Zr–10Nb alloy annealed at 600 ◦C. The inset is theenlarged part of the austenitic transformation during heating.

Fig. 6. (a) Stress–strain curve measured for the Ti–20Zr–10Nb alloy annealed at 600

Fig. 5. Stress–strain curves measured for the as-rolled Ti–20Zr–10Nb alloy (a), andfollowing annealed at different temperatures: (b) 600, (c) 700 and (d) 800 ◦C, respec-tively.

to �′′-phase martensite leads to a weak DSC peak upon cooling [9].The transformation characteristics of Ti–Zr–Nb annealed at othertemperatures are the same as that annealed at 600 ◦C.

Fig. 5 shows the tensile stress–strain curves of Ti–20Zr–10Nballoys tested at room temperature. It is seen from Fig. 5 that theas-rolled Ti–20Zr–10Nb alloy fractures after exhibiting only smallplastic deformation and the highest critical stress compared withthe annealed samples, which results from the work-hardeningeffect induced by cold rolling. For the annealed Ti–20Zr–10Nb alloysamples, the fracture strain decreases with the increasing anneal-ing temperature. As mentioned above, the Ti–20Zr–10Nb alloyannealed at 600 ◦C obtains the finest crystal grain compared withother samples annealed at higher temperature, and it is well knownthat grain refinement is beneficial to the plasticity of alloys, e.g. inNiMnGa SMA [26]. Each of the tensile stress–strain curves for theas-annealed Ti–20Zr–10Nb alloy consists of three obvious stageswhich are associated with the elastic deformation, reorientation ofmartensitic variants, corresponding to the stress plateau at about220–250 MPa, and the elastic/plastic deformation of reorientedmartensite. The tensile strength and failure strain of Ti–20Zr–10Nb

annealed at 600 ◦C are measured to be 542 MPa and 13.1%, respec-tively, superior to 200 MPa and 7% of the Ti–Nb–Pd alloy [13].

Fig. 6a shows the loading–unloading curves of theTi–20Zr–10Nb alloy annealed at 600 ◦C to different pre-strain.

◦C and (b) shape recovery strain of the Ti–20Zr–10Nb alloy versus pre-strain.

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212.[28] Y.Q. Wang, Y.F. Zheng, W. Cai, L.C. Zhao, Scripta Materialia 40 (1999)

1327–1331.

56 Y. Cui et al. / Materials Science a

ach stress–strain curve is obtained at room temperature by aoading and unloading tensile test and shape memory recoverytrain is measured as shown in a dot line by following heating upo 600 ◦C. The shape memory behaviors of Ti–20Zr–10Nb alloy are

easured as a function of pre-strain as summarized in Fig. 6b.t can be seen that the maximum SME is up to 2.5% for 8% totalre-strain with the corresponding shape recovery ratio of 30%, andME performance degrades for further tensile deformation. Theaximum SME value of Ti–20Nb–10Zr alloy compares favorablyith 0.7% for NiAl [27], equivalent to 3% for TiNiHf [28] and TiNb

2,8], less than 4.2% for polycrystalline NiMnGa [26] and 5.5% foriNiPd [29], however. The relatively low shape recovery strainf Ti–20Zr–10Nb alloy compared with NiMnGa and TiNi alloysight result from the low ordering bcc-� parent. However, it is

elieved that the shape memory behaviors of TiZrNb alloy could bemproved by some proper thermo-mechanical treatments, which

ill be further investigated.

. Conclusion

In summary, the grain size grows up with the increasing anneal-ng temperature and the refined grains are obtained in alloynnealing at 600 ◦C. The as-annealed Ti–20Zr–10Nb alloys areainly composed of �′′ martensite phase and a small amount ofphase. The alloy exhibits a reverse transformation temperature

igher than 500 ◦C. The TiZrNb alloy annealing at 600 ◦C obtainhe maximum tensile plasticity of 13.1% due to the refined grainize, and the stress for reorientation of martensitic variants and theensile strength of Ti–20Zr–10Nb are about 220 MPa and 540 MPa,espectively. The maximum SME of 2.5% has been obtained for 8%re-strain with the corresponding shape recovery ratio of 30%.

cknowledgements

This work is supported by National Natural Science Foundationf China (NSFC), No. 50771007, and Beijing Nova Programme.

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