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Materials Science and Engineering A 441 (2006) 112–118

Effect of Zr and Sn on Young’s modulus and superelasticityof Ti–Nb-based alloys

Y.L. Hao ∗, S.J. Li, S.Y. Sun, R. YangShenyang National Laboratory for Material Science, Institute of Metal Research, Chinese Academy of Sciences,

72 Wenhua Road, Shenyang 110016, PR China

Received 12 April 2006; received in revised form 9 July 2006; accepted 14 September 2006

bstract

Quaternary Ti–(20–26)Nb–(2–8)Zr–(3.5–11.5)Sn (wt%) alloys were investigated to evaluate the effects of Zr and Sn on Young’s modulusnd superelasticity of Ti–Nb-based alloys. X-ray diffraction analysis showed that solution-treated alloys have � + �′′, � + �, �′′ + �, �′′, or �icrostructures. Zr and Sn increase the lattice parameters of the � phase; for orthorhombic �′′ matensite, they increase the lattice parameter a but

ecrease both b and c. The martensitic start temperature of the �′′ is depressed by Zr and Sn additions, whereas the formation of athermal � isependent on Zr and Sn contents. Differential scanning calorimetry (DSC) measurements show that 1 wt% of Nb, Zr or Sn addition decreases theartensitic start temperature by 17.6, 41.2 or 40.9 K, respectively, due to their negative effect on lattice parameter ratios of the martensite (c/a and

/a). Tensile tests were used to evaluate Young’s modulus and superelasticity of the solution-treated alloys. Of the studied alloys Ti–24Nb–4Zr–7.5Snith single � microstructure has the lowest Young’s modulus of 52 GPa and recoverable elastic strain of about 2% at room temperature after cyclic

train.2006 Elsevier B.V. All rights reserved.

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eywords: Young’s modulus; Superelasticity; Biomedical titanium alloy; Mart

. Introduction

Titanium and its alloys have been widely used as biomedicalaterials to replace disfunctioned hard tissue in human body

ue to their light weight, low elastic modulus, high strengthnd excellent biocompatibility and corrosion resistance. Duringhe past decade, both near �-type and �-type titanium alloysomprising only non-toxic and non-allergic elements have beeneveloped in an effort to match high strength with low moduluso as to further ease “stress shielding” problem [1,2]. The addi-ion of high concentration of � stabilizers to some of the newlyeveloped alloys resulted in �′′ martensite with orthorhombictructure under the condition of quenching [3–5]. The presencef �′′ martensite, however, reduces hardness and tensile andatigue strengths of the alloys [4–8]. To avoid impairment to

ervice life, therefore, alloys containing the �′′ martensite areenerally considered unsuitable for making implant devices forard tissue replacement.

∗ Corresponding author. Tel.: +86 24 2397 1961; fax: +86 24 2390 2021.E-mail address: ylhao@imr.ac.cn (Y.L. Hao).

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921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.09.051

c transformation; Alloying effect

On the other hand, the martensitic and reverse transforma-ions concerning the �′′ can be exploited to develop shape

emory titanium alloys for bio-functional applications, as pio-eered by Baker and by Duerig et al. [9–11]. Their resultsuggested that shape memory effect is quite sensitive to heat-ng rate and the maximum recovered strain is about 3% underhe condition of salt bath heating with heating rate higher than0 ◦C/s. Increasing concern over allergic and toxic effects ofi ions released from TiNi led to recent research efforts towardeveloping Ni-free shape memory or superelastic titanium alloys12–19].

In �-type titanium alloys, athermal � phase is generallyormed in as-quenched conditions [20]. The influence of the

phase on shape memory effect and superelasticity of tita-ium alloys, however, is not well understood. Moffat and Lar-alestier [21] investigated systematically the competing forma-ion between the �′′ martensite and the � phase in quenchedinary Ti–Nb alloys. Their results suggested that the � phase

recipitates would not favour the formation of �′′ martensitend would impair shape memory effect and super-elastic prop-rties of titanium alloys. Recent experimental findings [22,23] ofuper-elastic deformation in the absence of �′′ martensite is not

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Y.L. Hao et al. / Materials Science

onsistent with the above conclusion. Reversible � to � phaseransformation or dislocation-free deformation mechanism haveeen suggested as possible origins of the superelastic behaviour22,23].

The objective of this study is to investigate the effect of Nb, Zrnd Sn on phase formation, Young’s modulus and superelasticityn as-quenched titanium alloys, in the hope of reducing Young’s

odulus and improving super-elastic properties by alloying withr and Sn.

. Experimental

Quaternary Ti–Nb–Zr–Sn alloys were melted in an arc melt-ng furnace using a tungsten electrode under argon protectionith magnetic agitation. They were melted three times using ai–Sn master alloy and pure Ti, Nb and Zr as raw materials. Theominal chemical compositions of alloys are listed in Table 1all in wt%). The 40 g buttons obtained were forged at 950 ◦Co 10 mm diameter cylinders. These cylinders were encapsuledn evaluated quartz tubes with pressure of 10−2 Pa and wereolution-treated at 850 ◦C for 1 h before quenching into watery breaking the capsules. Two buttons with different compo-itions were analysed by wet chemical analysis to characterizehe differences between nominal and actual compositions, andhe results are shown in Table 2. Interstitial contents of N, Hnd O of both alloys were also listed in Table 2. The nominalompositions shall be used to denote the alloys hereafter.

Tensile testing was conducted at room temperature (21 ◦C) byyclic loading at initial strain rate of 1 × 10−3 s−1 using spec-mens with a gauge section of 3 mm in diameter and 15 mmong by MTS 810 Material Test System. In order to improvehe accuracy of measurement, tensile Young’s modulus andecovered strains were determined from the stress–strain curvesecorded by using a strain gauge. The Vickers hardness was

easured under a load of 10 kg applied for 15 s along longitudi-

al direction of mechanically ground and polished specimens.ransformation temperature was measured by differential scan-ing calorimetry (DSC) at heating or cooling rate of 10 ◦C/min

able 1ominal chemical compositions (wt%) and corresponding phase constitutionsf as quenched Ti–Nb–Zr–Sn alloys determined by X-ray diffraction analysis

20Nb 22Nb 24Nb 26Nb

Zr–7.5Sn � + �′′Zr–7.5Sn � + �′′ � + �′′ � �

Zr–7.5Sn � + �

Zr–3.5Sn �′′ + � �′′ + � � + �′′ � + �′′Zr–11.5Sn � + � � + � � + �

able 2hemical analysis of Ti–Nb–Zr–Sn alloys (wt%)

omposition Nb Zr Sn O N H

ominal 24 4 7.5nalysed 24.0 3.93 7.54 0.069 0.0076 0.0049ominal 26 4 7.5nalysed 26.1 3.89 7.35 0.063 0.0079 0.0044

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ngineering A 441 (2006) 112–118 113

sing Perkin-Elmer Pyris Diamond Type calorimeter. Phase con-titutions and their lattice parameters in as-quenched specimensere determined by 2θ/θ coupling method of X-ray diffraction

nalysis along the longitudinal direction of specimens using/max 2500PC Rigaku diffractometer. In order to avoid arti-

act of stress-induced martensitic transformation, the specimensere heavily etched in a water solution with 8 vol.% of HF to

emove surface layer with internal stress introduced by grindingnd polishing. In order to increase accuracy of lattice parame-er measurement, a low scanning speed of 1◦/min was adoptedn this study. Specimens for optical microscope observationere etched at the boiling temperature of a water solution with0 vol.% HCl. Transmission electron microscopy (TEM) speci-ens were prepared from mechanically-thinned plates by elec-

ropolishing in a solution of 21% perchloric acid, 50% methanolnd 29% n-butyl alcohol at about −40 ◦C. The thin foils werexamined on a Philips EM420 transmission electron microscopeperating at 100 kV.

. Results and discussions

.1. Microstructures and phase constitutions ofi–Nb–Zr–Sn alloys

Fig. 1 shows three typical microstructures (opticalicrographs) of as-quenched quaternary alloys, represented

y Ti–24Nb–4Zr–7.5Sn, Ti–24Nb–4Zr–3.5Sn and Ti–20Nb–Zr–7.5Sn alloys. They appear to have single � microstructureFig. 1(a)), or � matrix with high and low amount of the �′′artensite (Fig. 1(b and c)). The small dark spots in Fig. 1(a) are

ften observed on metallographic samples of heavily deformed-phase titanium alloys and are probably dislocation etch pits.hese metallographs also show that the average grain size of thephase is about 80 �m.Systematic X-ray diffraction analyses revealed that alloys of

he studied range of composition have � + �, �′′ + �, � + �′′,′′, or � microstructures (Table 1). It is clear from Table 1 that,

or identical Zr and Sn contents, the �′′ martensite is depressedith the increase of Nb content. This is consistent with exper-

mental results for binary Ti–Nb alloys [21]. A similar trendas also observed in the effects of Zr and Sn on the martensitic

ransformation, as demonstrated by X-ray diffraction profiles ofi–24Nb-based alloys presented in Fig. 2: under the conditionf 7.5 Sn addition, the �′′ martensite forms in the alloy with 2%r (Fig. 2(e)) but is suppressed in the alloys with 4 and 8% Zr

Fig. 2(c and d)). For 4Zr addition, the �′′ martensite appears inlloy with low Sn addition (Fig. 2(b)) but disappears with thencrease of Sn contents (Fig. 2(a and d)). Clearly, the additionf Nb, Zr and Sn decreases the stability of the �′′ martensite, asonfirmed by the variation of lattice parameters with chemicalompositions to be described in the next section.

The effects of Zr and Sn on � phase formation are com-licated. Data presented in Table 1 and Fig. 2 show that, for

i–24Nb–7.5Sn base composition, the � phase appears ifr addition is 8% but disappears for lower Zr contents (2%nd 4%); for Ti–24Nb–4Zr base composition, the � phase isuppressed for 3.5% and 7.5% Sn addition but forms if Sn

114 Y.L. Hao et al. / Materials Science and Engineering A 441 (2006) 112–118

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ig. 1. Typical optical microstructures of water quenched quaternaryi–Nb–Zr–Sn alloys: (a) Ti–24Nb–4Zr–7.5Sn, (b) Ti–24Nb–4Zr–3.5Sn and (c)i–20Nb–4Zr–7.5Sn alloys.

ontent is increased to 11.5%. By contrast, in Ti–20Nb–4Zr andi–22Nb–4Zr base compositions, the � phase appears for Sndditions of 3.5% and 11.5% but disappears for intermediatemount of Sn (e.g., 7.5%). Clearly, the effects of Zr and/or

n on the � to � transformation is complex compared to theironotonic influence on the � to �′′ transformation.Competition between the � phase and the �′′ martensite

uring quenching has been investigated in metastable � type

α

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ig. 2. X-ray diffraction profiles of alloys based on Ti–24Nb with: (a)Zr–11.5Sn, (b) 4Zr–3.5Sn, (c) 8Zr–7.5Sn, (d) 4Zr–7.5Sn and (e) 2Zr–7.5Sndditions.

itanium alloys. Moffat and Larbalestier [21] conducted trans-ission electron microscopy (TEM) analysis of alloys with

ifferent Nb contents and concluded that the modes of �hase decomposition to both the � phase and the �′′ marten-ite are mutually exclusive. The � phase forms generally inetastable � type titanium alloys if the formation of the �′′artensite is suppressed during water quenching from above

he � transus [24]. In the present investigation, TEM analysisFig. 3) did not find characteristic diffraction spots due to the �hase in Ti–24Nb–4Zr–7.5Sn alloy that features single � phaseicrostructure (Figs. 1(a) and 2(d)). This suggests that appro-

riate amounts of Zr and Sn have the advantage of suppressinghe formation of both the martensite and the � phase when thelloy is cooled to room temperature.

The absence of (1 1 1) and (0 1 2) superlattice diffractioneaks of the � phase in X-ray diffraction profiles of alltudied alloys suggests that the � phase is disordered. Thiss consistent with the results for Ti–29Nb–13Ta–4.6Zr andi–39Nb–13Ta–4.6Zr alloys examined by TEM [24]. By con-

rast, Banerjee et al. [25] recently reported that the metastable �hase in Ti–34Nb–9Zr–8Ta alloy is ordered.

.2. Effects of Zr and Sn on lattice parameters of the β and

′′

Lattice parameters of both the � phase and the �′′ marten-ite determined by X-ray diffraction analysis are listed in

Y.L. Hao et al. / Materials Science and Engineering A 441 (2006) 112–118 115

Fig. 3. Bright field TEM micrograph (a) and {1 1 0}� selected area diffractionpattern (b) of Ti–24Nb–4Zr–7.5Sn alloy.

Table 3Lattice parameter (A) of � phase in Ti–Nb–Zr–Sn alloys

20Nb 22Nb 24Nb 26Nb

2Zr–7.5Sn 3.2914Zr–7.5Sn 3.295 3.299 3.3048Zr–7.5Sn 3.30844

Tate

TL

2222222

Table 5Lattice strains of � to �′′ transformation in Ti–Nb–Zr–Sn alloys

〈−1 1 1〉�/〈0 0 1〉�′′ (%)

〈0 0 1〉�/〈0 1 0〉�′′ (%)

〈1 1 0〉�/〈1 0 0〉�′′ (%)

24Nb-2Zr–7.5Sn 1.1 4.6 −5.222Nb–4Zr–7.5Sn 0.64 5.0 −5.226Nb–4Zr–3.5Sn 0.35 5.9 −6.024Nb–4Zr–7.5Sna 0.16 3.8 −4.3

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Zr–3.5Sn 3.293Zr–11.5Sn 3.297 3.299 3.302

ables 3 and 4, respectively. In agreement with the binary Ti–Nblloys [21], the addition of Nb increases the lattice parameter ofhe � phase. The data in Table 3 also show that the lattice param-ter increases with Zr up to 8% content, whereas its increase with

able 4attice parameters (A) of the orthorhombic �′′ martensite in Ti–Nb–Zr–Sn alloys

a b c c/a b/a

4Nb-2Zr–7.5Sn 3.119 4.868 4.703 1.508 1.5610Nb–4Zr–7.5Sn 3.089 4.942 4.709 1.524 1.6002Nb–4Zr–7.5Sn 3.123 4.892 4.691 1.502 1.5680Nb–4Zr–3.5Sn 3.050 4.981 4.702 1.542 1.6332Nb–4Zr–3.5Sn 3.074 4.947 4.692 1.526 1.6094Nb–4Zr–3.5Sn 3.096 4.928 4.674 1.510 1.5926Nb–4Zr–3.5Sn 3.134 4.888 4.664 1.488 1.560

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a The lattice parameters of the �′′ martensite of this alloy are extrapolatedrom Ti–(20, 22)Nb–4Zr–7.5Sn alloys.

n is obvious up to 7.5% but then remains almost constant withurther increase up to 11.5%.

As to the �′′ martensite, Table 4 shows that Nb tends toncrease the a but decrease b and c of the orthorhombic lat-ice; the lattice parameter ratios of b/a and c/a decrease withhe increase of Nb content as a result. Similar tendency is alsobserved for Zr and Sn from Table 4. For binary Ti–Nb alloys,revious investigations showed that phase transformation startemperature of the �′′ martensite decreases with the decreasef lattice parameter ratios of b/a and c/a, as reviewed in [21].hus, the depressing effects of Zr and Sn on the �′′ martensite

ormation can be related to their effects on lattice parameteratios.

Lattice strains accompanying the � to �′′ phase transfor-ation in quaternary Ti–Nb–Zr–Sn alloys are estimated and

resented in Table 5. A dependence on chemical compositions clearly seen and the minimum lattice strains occur at chemi-al compositions close to Ti–24Nb–4Zr–7.5Sn. This quaternarylloy has much reduced lattice strains compared with ternaryi–Nb–Sn alloys reported in [25].

.3. Transformation temperatures of Ti–Nb–Zr–Sn alloys

The temperatures for the � to �′′ and reverse phase trans-ormations in quaternary alloys were measured by DSC methodetween ±150 ◦C at cooling and heating rates of 10 ◦C/min. Inharp contrast with NiTi shape memory alloy, phase transfor-ation peaks of the studied alloys are quite weak; only those

n several compositions can be distinguished and the results arehown in Table 6. The peak heights of the phase transformationn the alloys listed in Table 6 are generally lower than 0.02 W/gnd even difficult to distinguish. Fig. 4 gives an example of DSCeasurement of Ti–24Nb–4Zr–7.5Sn alloy. The peak heights of

he studied alloys are also lower than those reported for ternaryi–Nb–Sn alloys [13].

In order to examine the origin of weak phase transforma-ion peaks of DSC measurements, polished Ti–24Nb–4Zr–7.5Sn

able 6emperatures of �–�′′ phase transformation in Ti–Nb–Zr–Sn alloys

Ms (◦C) Mf (◦C) As (◦C) Af (◦C)

6Nb–4Zr–7.5Sn −77.5 −78.84Nb–4Zr–7.5Sn −50.1 −53.2 −44.2 −42.44Nb–4Zr–3.5Sn 113.7 112.54Nb–2Zr–7.5Sn 32.3 30.2

116 Y.L. Hao et al. / Materials Science and E

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Fig. 4. DSC curves of water quenched Ti–24Nb–4Zr–7.5Sn alloy.

pecimens were subjected to in situ optical microstructure obser-ation during cooling and heating between 20 and −150 ◦C.he results show that the transformation from the � matrix to′′ martensite is quite limited when the specimen is cooled

o −80 ◦C (Fig. 5), ∼30 ◦C lower than the martensitic fin-sh temperature measured by DSC (Table 6). This suggestshat the martensitic transformation is not fully thermoelastic

ig. 5. In situ optical microstructure observation of Ti–24Nb–4Zr–7.5Sn alloyater quenched to 20 ◦C (a) and −80 ◦C (b).

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ngineering A 441 (2006) 112–118

n the studied Ti–Nb–Zr–Sn quaternary alloys, in contrast withully thermoelastic transformations during cooling and heatingbserved in Ti–Nb–Sn ternary alloys [13]. The in situ opticalicrostructure observation also shows that the transformation

tart temperature is about −10 ◦C, higher than the DSC mea-ured value of about −50 ◦C, suggesting internal stress inducedy surface polishing has significantly increased the martensiticransformation temperature.

Comparing the data of Ms temperature shown in Table 6,t can be estimated that 1 wt% of Nb, Zr or Sn decreases

s temperature by about 17.6, 41.2 or 40.9 ◦C, respectively.he estimation was made using actual chemical compositionsbtained by chemical analysis instead of nominal ones listedn Table 1. It also should be noted that the influence of inter-titial elements such as oxygen was assumed constant in thebove calculations. Based on these data, Ms temperatures of thelloys not listed in Table 6 can be roughly estimated: the phaseransformation temperatures of Ti–(20–26)Nb–4Zr–11.5Sn andi–24Nb–8Zr–7.5Sn alloys are lower than −150 ◦C, while thosef Ti–(20, 22)Nb–4Zr–3.5Sn alloys are higher than +150 ◦C.hey are all out of the temperature range we set for DSC study.

The effects of Nb and Sn on martensitic start (Ms) tem-erature have been investigated in ternary Ti–Nb–Sn alloys byanada and coworkers [13,26]. Inconsistent results, with 1% Snecreasing Ms by about 78 ◦C [13] or 52 ◦C [26], were reported.oth reported data were much higher than the estimation of theresent study (40.9 ◦C). In addition, the Ms decrease due to Nbn ternary alloys [26] is about twice that in the studied quaternarylloys. One possible reason of the above quantitative differences related to chemical composition. The following experimen-al results tend to support such speculation: Sn was reported toave quite weak effect on martensitic transformation in binaryi–Sn but the effect is strong in ternary Ti–Nb–Sn alloys [26];dditional evidence is provided by small addition of Pd in Ti–Nblloys that weakens the effect of Nb on martensitic transforma-ion and results in trivial change of transformation temperatureven when Nb content is increased from 30 to 40 wt% [19]. Howhe interaction of these alloying elements influences the marten-itic transformation temperature is open for further theoreticalnvestigation.

Table 6 shows that the differences of Ms and Mf as well ass and Af are just about 2 ◦C, much lower than those for ternaryi–Nb–Sn alloys [13,26]. Furthermore, the difference betweenf and Ms is about 7.7 ◦C in Ti–24Nb–4Zr–7.5Sn alloy, much

ower than Ti–Nb–Sn ternary alloys. Takahashi et al. [13] anditta et al. [26] reported that the corresponding difference is

bout 25 and 90 ◦C in Ti–25.4Nb–9.9Sn and Ti–27Nb–8.1Snlloys, respectively. This small difference in transformation tem-eratures in Ti–24Nb–4Zr–7.5Sn alloy can be attributed to theow lattice strains of � to �′′ phase transformation, as men-ioned in Section 3.2. It was noted earlier that the martensiticnd reversible transformations are difficult to complete in theuaternary alloys during cooling and heating (see Fig. 5). This is

onsistent with the small differences of start and finish transfor-ation temperatures, Ms and Mf as well as As and Af. However,

t cannot explain the small difference between both start tem-eratures, As and Ms.

Y.L. Hao et al. / Materials Science and E

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etim�afsnot go to completion. The double yielding phenomenon is alsoobserved for Ti–24Nb–4Zr–7.5Sn alloy (Fig. 7(b)) but not forTi–24Nb–4Zr–11.5Sn alloy (Fig. 7(c)). In this case, the � phaseis fully stabilised with respect to the martensitic transformation

ig. 6. Tensile Young’s modulus (a) and Vickers hardness (b) of water quenchedi–Nb–Zr–Sn alloys.

.4. Young’s modulus and superelasticity of Ti–Nb–Zr–Snlloys

Tensile Young’s moduli of the studied alloys are plotted inig. 6(a). The Young’s modulus of as-quenched alloys variesith chemical composition and a minimum of about 52 GPa isbtained in Ti–24Nb–4Zr–7.5Sn alloy. Examining phase consti-utions of the alloys given in Table 1, it can be seen that the alloysontaining the � phase generally have high Young’s modulus.his is because the � phase has higher Young’s modulus than

he � and the �′′ phases [4–6]. The data in Fig. 6(a) also showhat Young’s moduli of the alloys containing the � phase arenly slightly higher than those without the � phase. This sug-ests that the volume fractions of the � phase are quite low inhese alloys. The measurements of Vickers hardness presentedn Fig. 6(b) also support this deduction.

Superelasticity of water quenched Ti–Nb–Zr–Sn alloysas evaluated by cyclic deformation at initial strain ratef 1 × 10−3 s−1 at room temperature (21 ◦C). Three typi-al stress–strain curves, demonstrated by Ti–24Nb-2Zr–7.5Sn,i–24Nb–4Zr–7.5Sn and Ti–24Nb–4Zr–11.5Sn alloys, are

hown in Fig. 7. The martensitic start (Ms) temperature forhe above three alloys decreases in turn. Although the tensileest temperature is about 9 ◦C lower than the martensite finishMf) temperature of Ti–24Nb–2Zr–7.5Sn (Table 6), the � par-

F(

ngineering A 441 (2006) 112–118 117

nt phase did not fully transform to the �′′ martensite beforeesting. This is evidenced by the double yielding phenomenonn the stress–strain curve (Fig. 7(a)) due to the stress-assisted

artensitic transformation. X-ray analysis (Table 1) that reveals+ �′′ microstructure of Ti–24Nb–2Zr–7.5Sn supports above

rgument. As mentioned in Section 3.3, the martensitic trans-ormation in the studied quaternary alloys appears much moreluggish than is usual for thermoelastic martensite and does

ig. 7. Stress–strain curves at 21 ◦C of water quenched Ti–24Nb–2Zr–7.5Sna), Ti–24Nb–4Zr–7.5Sn (b) and Ti–24Nb–4Zr–11.5Sn (c).

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ecause its Ms temperature is too low as a result of its high Snontent.

Since elastic strain recorded by a tensile testing machine cane twice as high as its true value [13], the stress–strain curveshown in Fig. 7 were recorded by using a strain gauge to ensureccurate measurement of elastic strain. It can be seen fromig. 7(b) that, for Ti–24Nb–4Zr–7.5Sn alloy, elastic strains of%, 2% and 3% are totally, near totally and partially recovered,espectively, after unloading. The superelasticity of this alloy islightly lower than ternary Ti–25.4Nb–9.9Sn alloy [13]. Fromhe view point of thermoelastic transformation, the above dif-erence can be explained by the difference between the testingemperature and the austenitic finish (Af) temperature. For theernary alloy, the testing temperature is 30 ◦C higher than thef temperature, a generally accepted temperature difference forchieving significant superelasticity. On the other hand, the tem-erature difference for the quaternary alloy, about 63 ◦C, is tooigh to obtain good superelasticity. The above argument suggestshat increasing Af temperature by decreasing the contents ofb, Zr and/or Sn would improve superelasticity. However, addi-

ional experiments yield contrary results. Quaternary alloys withncreased Sn content were found to possess improved supere-asticity and a maximum recoverable tensile strain of 3.3% wasbtained in an alloy with 7.9 wt% Sn (compared to 7.5 wt% Snn the studied alloy) [27]. Furthermore, stress–strain curve ofs hot-rolled 7.9Sn alloy exhibits non-linear elastic deformationehaviour without double yielding, substantially different fromhe curve shown in Fig. 7(b).

It should be noted that there exist several reports of supere-asticity that are not related to the �′′ martensitic transforma-ion [19,22,23]. Further investigation to reveal other kinds of

echanism is therefore of crucial importance for improving theuperelastic properties of titanium alloys (e.g., by combiningwo or more mechanisms).

. Conclusions

The effects of Zr and Sn contents on Young’s modulusnd superelasticity of water quenched quaternary Ti–Nb–Zr–Snlloys were investigated in this study. The main results are sum-arized below:

1) Both Zr and Sn tend to increase the lattice parameter of the� phase; they increase the lattice parameter a but decreaseb and c of the �′′ martensite.

2) The martensitic start temperature decreases by about 17.6,41.2 and 40.9 ◦C due to the addition of 1 wt% of Nb, Zrand Sn, respectively. Zr and Sn are therefore effective insuppressing the �′′ martensitic transformation, as can beexplained by their effects on the lattice parameter ratios of

c/a and b/a.

3) The lattice strains of forming the orthorhombic �′′ marten-site from the bcc � phase are dependent on chemical com-position. The minimum is obtained in Ti–24Nb–4Zr–7.5Sn

[

ngineering A 441 (2006) 112–118

among the studied alloys. A small temperature hysteresis,Af–Ms, of about 7.7 ◦C, is detected in this alloy.

4) The formation of the athermal � is dependent on both Zrand Sn contents in Ti–(20–26)Nb-based alloys.

5) Ti–24Nb–4Zr–7.5Sn alloy with single � microstructure hasa minimum Young’s modulus of 52 GPa and recovered elas-tic strain larger than 2% after cyclic deformation at roomtemperature.

cknowledgement

The work was partly supported by the NSFC (grants0471074 and 30471754) and the Chinese MoST (grantG2000067105).

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