Nb substitution on superelastic behavior of Ti–Nb–Zr alloy

Materials Science & Engineering A 563 (2013) 78–85

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Influence of equiatomic Zr/Nb substitution on superelastic behavior of Ti–Nb–Zr alloy Jinyong Zhang a, Fan Sun a, Yulin Hao b, Nicolas Gozdecki a, Emilie Lebrun a, Philippe Vermaut a, Richard Portier a, Thierry Gloriant c, Pascal Laheurte d, Fre´de´ric Prima a,n a

Laboratoire de Physico-Chimie des Surfaces, Groupe de Me´tallurgie Structurale, UMR 7045, ENSCP, 11 rue Pierre et Marie Curie, F-75231 Paris cedex 05, France Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China c ´tallurgie, 20, Avenue des Buttes de Coesmes, 35043 Rennes cedex, France Institut des Sciences Chimiques de Rennes, UMR CNRS 6226 Sciences Chimiques de Rennes/Chimie- Me d Laboratoire LEM3 UMR CNRS 7239, universite´ de Metz, France b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 15 August 2012 Received in revised form 1 October 2012 Accepted 14 November 2012 Available online 19 November 2012

Based on a binary Ti–26Nb (at%) alloy, Ti–(26-z)at% Nb-(z)at% Zr (z¼ 2, 6, 8 and 10) alloys via equiatomic substitution of Nb by Zr are formulated. Influence of equiatomic Zr/Nb substitution on microstructure evolution, mechanical properties and deformation mechanism of the superelastic Ti–Nb–Zr alloy are investigated. Experimental results show that the phase constitution is single b phase or b/oath phase at 0 o Zr/Nb o 0.35 since the measured Ms temperature is maintained at around 250 K. The two-phase microstructure consist of a00 martensite and b phase is obtained at Zr/Nb ratio 4 0.4 due to the attenuation of b stabilizing effect of Zr, which is related closely to the Nb content in ternary Ti–Nb–Zr alloys. The mechanism of the superelastic behavior alters gradually from reversible b/a00 martensitic transformation to rearrangement of pre-existing a00 martensites as a function of Zr/Nb ratio increase. At Zr/Nb ¼ 0.3, the alloy of corresponding composition exhibits the best superelasticity and combined mechanical performance. A coefficient of DT (DT ¼Tb-Ms) is proposed to understand the experimental results by evaluating the b instability of Ti–Nb–Zr alloys. & 2012 Elsevier B.V. All rights reserved.

Keywords: Microstructure characterization Mechanical properties Superelasticity Metastable b titanium alloys Equiatomic substitution

1. Introduction Titanium and its alloys are widely used for hard tissue replacement, cardiac and cardiovascular devices, due to their unique combination of mechanical properties such as high strength and fatigue resistance, low modulus, good ductility, good wear resistance and superior biocompatibility. During the last few decades, the Ni-free metastable b-type titanium alloys exhibiting superelasticity (SE) and shape memory effect (SME) such as Ti–V-based alloys [1,2], Ti–Mo-based alloys [3–10] and Ti–Nbbased alloys [11–24] have attracted increasing interest in biomedical applications. However, Ti–V-based alloys are not suitable for biomaterial because of the cyto-toxity of V, Ti–Mo-based alloys and Ti–Nb-based alloys seem to be better candidates in developing metastable b-type biomaterial. The shape memory effect (SME) and superelasticity (SE) are associated with the reversion of b to a00 in b-type Ti-based alloys. Actually, systematic researches which characterize the reversible martensitic transformation (MT) between the b parent phase and the a00 martensite and its influence on the mechanical properties, shape

n

Corresponding author. Tel.: þ33 1 43 54 87 02, fax: þ33 1 44 27 67 10. E-mail address: [email protected] (F. Prima).

0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.11.045

memory effect (SME) and superelastic (SE) deformation behavior on Ti–Nb binary alloys and Ti–Nb-based alloys, have been undertaken by several groups [25,26]. The results show that the shape memory effect and superelasticity in Ti–Nb alloys can be improved by adding other alloying elements such as Ta [27,28], Zr [12,27], Sn [29,30], Al [31] and/or interstitial elements such as O and N [27,32], and large recoverable strains can be also obtained from the stress-induced martensitic transformation (MT). These achievements have brought a better understanding of the reversible martensitic transformation (MT) in that category of materials for further development of new improved superelastic alloys. Zr addition in Ti–Nb alloys has been shown to improve the SME and the superelasticity by the increase of phase transformation strains [12,33,34]. The Ti–Nb–Zr ternary alloys therefore have a promising potential for biomedical applications. An interesting point has been noted that Zr element is thought as a kind of neutral elements due to their weak effect on b transus temperature but depress significantly the Ms temperature as combined with large amount of Nb [29,35] in Ti–Nb alloys. It has been noted that Zr and Nb elements have previously shown to display similar effect on decreasing Ms temperature [12,16,36]. For example, the Ms temperature decreases by 38 K per 1 at% increase of Zr content for Ti–22Nb–(2–8)Zr(at%) and by 40 K per 1 at% increase of Nb content for Ti–(20–28)Nb(at%). However,

J.Y. Zhang et al. / Materials Science & Engineering A 563 (2013) 78–85

they play a different role in affecting the b-transus temperature (Tb), where Zr is known as a neutral element with little influence on Tb whereas Nb behaves as a classical b stabilizer in decreasing Tb temperature. As a consequence, equiatomic substitution of Nb by Zr could result in an increase of Tb temperature by keeping the Ms temperature at almost the same level. For instance, our previous investigation [37] showed that Tb temperature was increased by 70 K in the Ti–20Nb–6Zr (at%) compared to that of Ti–26Nb(at%) while Ms temperature was kept almost the same level. The subsequent change of b phase stability, induced by the equiatomic substitution, is actually promoted the superelasticity at room temperature due to its positive influences on material’s stacking fault energy (SFE) and the reversible b/a00 martensitic transformation. In this paper, following the Ti–26Nb alloy, the Ti–(26-z)at% Nb–(z)at% Zr (z¼2, 6, 8 and 10) alloys via equiatomic substitution of Nb by Zr were formulated. We focused on the effect of equiatomic Zr/Nb subsititution on microstructure evolution, mechanical properties and deformation mechanism of the superelastic Ti–Nb–Zr alloy. The Tb temperature of the investigated alloys are estimated using an empirical method [38]: Tb ¼1158 8.5(Nb)–2(Zr), in which the element in square brackets describes its wt% concentration in the present alloy. Based on the Ms temperature of Ti–26Nb ( 250 K), the Ms temperature of the investigated alloys via equiatomic substitution of Nb by Zr are estimated by the equation: Ms ¼250þ40z 38z

Table 1 Calculated Tb, estimated Ms, e/a ratio and Mo equivalent of each composition. Alloys

Tb

Ms

e/a

Mon

Ti–16Nb–10Zr Ti–18Nb–8Zr Ti–20Nb–6Zr Ti–22Nb–4Zr Ti–24Nb–2Zr Ti–26Nb

914 894 874 853 833 813

270 266 262 258 254 250

4.16 4.18 4.20 4.22 4.24 4.26

7.00 7.87 8.74 9.61 10.48 11.35

n

[Mo] Eq. [39]¼Mo þ Nb/3.6

79

(Ms decrease by 40 K per 1 at% for Nb element, by 38 K per 1 at% for Zr element), where z¼2, 4, 6, 8 and 10. Meanwhile d-electron design method developed by Morinaga which aims to design the superelastic Ti-alloys by concerning the physical background in phase stability and phase transformation, is also considered in this study. However, the particular effect of Zr (nearly neutral on Tb) in b stability could lead to inaccuracy of d-electron design method, another criteria, the e/a ratio and equivalent Mo content [39], are also taken into account in characterizing the compositions. Table 1 lists the parameters of calculated Tb, estimated Ms, e/a ratio and equivalent Mo content of each composition. The investigations on their mechanical properties and microstructure are analyzed in order to clarify the influences of Zr on martensitic transformation and b stability by substituting Nb in ternary Ti–Nb–Zr alloy.

2. Experimental Five alloys of Ti–(26-z) at% Nb–(z) at% Zr (z ¼0, 2, 6, 8 and 10) (all compositions hereafter are described in atomic percent) were synthesized by Ar arc-melting method using high-purity titanium, niobium and zirconium. In the melting process, the alloy ingots about 25 g were melted for five times, and flipped over each time before melting, then finally formed into cylinder shape. Ingots obtained from melting were homogenized at 1173 K for 72 ks under high vacuum of 10  6 Pa, followed by water quenching. The as-quenched ingots were heavily cold rolled to 0.5 mm in thickness at a reduction rate of more than 95% at room temperature. From the cold-rolled sheet, tensile specimens were mechanically prepared with gage dimensions of 50 mm  5 mm  0.5 mm. The specimens were solution-treated (ST) in the b-phase domain at 1173 K for 1.8 ks under high vacuum of 10  6 Pa in order to restore a fully recrystallized microstructure. The phase constitution of the samples were characterized by X-ray diffraction (XRD) using PANalytical X’Pert Pro diffractometer with CuKa radiation operating at 45 kV and 40 mA. The differential scanning calorimetric (DSC) measurement, in order to determine the Ms

Fig. 1. Optical micrographs of the solution-treated Ti–24Nb–2Zr (a), Ti–20Nb–6Zr (b), Ti–18Nb–8Zr (c) and Ti–16Nb–10Zr (d) alloys, respectively.

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temperature, was applied by using METTLER TOLEDO DSC822e equipment at the rate of 20 K/min. Microstructures of the samples were observed by optical microscopy (OM) and transmission electron microscopy (TEM). Specimens for OM were first mechanically polished on silicon carbide abrasive papers followed by a final polishing step with a colloidal silica suspension (particle size:50 nm) and next chemically etched with a solution H2O, HNO3 and HF (5:3:1) (vol%). The TEM investigations were performed by using a JEOL 2000FX microscope operating at 200 kV. Specimens for TEM observations were prepared by a conventional twin-jet polishing technique. The mechanical properties were evaluated by using an INSTRON 5966 machine with a strain rate of 10  3 s  1. An extensometer was used to precisely measure the deformation of the specimens. The superelastic behavior was characterized from cyclic loading–unloading tensile tests with strain increments of 0.5% till a total elongation of 5%. All tensile tests were performed along the rolling direction.

3. Results 3.1. Microstructure and phase constitution

precipitated during the quench process [40]. The a00 phase could not be detected on the ST Ti–24Nb–2Zr and Ti–20Nb–6Zr alloys due to the martensite start temperatures (Ms), around 240 K, are lower than the room temperature. On the other hand, the coexistence of a00 martensite and b phase in the ST Ti–18Nb–8Zr and Ti–16Nb–10Zr alloys are detected (shown in Fig. 2(c) and (d)) due to the actual martensite start temperature (Ms) are close to or above room temperature. Compared with the Ti–18Nb–8Zr alloy, a00 martensite seems to compose the majority of the microstructure in the Ti–16Nb–10Zr alloy. Meanwhile, the intensity of the (2 0 0) and (2 1 1) beta peak for Ti–18Nb–8Zr alloys as well as the (1 1 0) beta peak for Ti–20Nb–6Zr are much weaker than those of the corresponding beta peaks for other alloys. The phenomenon is quite common in Ti-based alloys from literatures [8,12,14,19,20] due to various reasons including testing method, shape effect of the specimen, secondary phases and microstructural texture from cold deformation and recrystallization. The qualitative XRD analysis of the specimens shows evolutional combinations of beta phase, omega phase and a00 martensite phase at the ST state. However, such phase constitution analysis with relatively low specific intensities provides limited information to clarify the microstructural texture at the ST state.

The microstructural characterization of ST specimens is analyzed using optical microscopy and TEM investigation. Fig. 1 shows optical micrographs of the four as-quenched alloys investigated in this study. The Ti–24Nb–2Zr and Ti–20Nb–6Zr (Fig. 1(a) and (b)) alloys exhibit a single equiaxed b-grain microstructure with grain sizes varying between a few tens and a few hundreds of micrometers. The precipitation of nanometric athermal omega (oath) phase cannot be detected with optical microscopy due to the small size of these particles. The Ti–18Nb– 8Zr and Ti–16Nb–10Zr (Fig. 1(c) and (d)) alloys present a dual phase microstructure, composed by equiaxed b grains with intragranular thin a00 needles. It can be noticed that the volume fraction of a00 in the Ti–16Nb–10Zr alloy is larger than that in Ti–18Nb–8Zr alloy. Fig. 2 shows the X-ray diffraction pattern of the four as-quenched alloys samples. The XRD pattern obtained on the as-quenched Ti–24Nb–2Zr shows diffractions peaks corresponds to single b phase (Fig. 2(a)). It has been reported that the metastable b state can be obtained in solution treated Ti–26Nb (at%) [37] and Ti–22Nb–4Zr (at%) [12] alloys at room temperature. In the ST Ti–20Nb–6Zr specimen, minor oath peaks in b matrix are detected from XRD (Fig. 2(b)). It is consistent with previous results that a very high density of athermal o nano-particles

XRD profiles on ST states

Intensity(a.u.)

110β 020α'' 111α'' 110α''

021α'' 110β

112α'' 200β 200α''130α'' 211β 113α'' (d) 112α'' 200α'' 130α'' 211β113α''

(c)

200β 0002ω

0001ω 110β 30

40

200β 50

60

211β

(b) 211β

(a) 70

80

2θ (degree) Fig. 2. X-ray diffraction profiles of the solution-treated Ti–24Nb–2Zr (a), Ti–20Nb–6Zr (b), Ti–18Nb–8Zr (c) and Ti–16Nb–10Zr (d) alloys, respectively.

Fig. 3. TEM images of the solution-treated Ti–20Nb–6Zr specimen, (a) bright-field image of the b grains in size of 20–30 mm, (b) diffraction pattern and corresponding dark field image of a single variant of nano o precipitates in large density and in small size (o 10 nm).

J.Y. Zhang et al. / Materials Science & Engineering A 563 (2013) 78–85

Fig. 3 shows a typical bright field TEM image (a) and the corresponding electron diffraction pattern (b) of the ST Ti–20Nb– 6Zr alloy. The fully recrystallized b grains are in the sizes of around 20 um to 30 um (Fig. 3(a)). The selected area electron diffraction shows typical diffuse diffraction spots of two variants of oath precipitates (inset of Fig. 3(a)). The corresponding dark field image indicates a high volume fraction of oath phase (shown in Fig. 3(b)), which confirms the results obtained by XRD analysis (Fig. 2(b)). However, for the ST Ti–16Nb–10Zr specimen, high volume fraction of a00 martensite was observed by TEM. Fig. 4 shows a typical bright field image of the microstructure in the ST Ti–16Nb–10Zr specimens, where complex morphology of a00 martensites and b phase is composed in different regions all over the image. The selected area electron diffraction pattern (Fig. 4(b)) and the corresponding dark field image illustrate two a00 variants in twinning relationship (Fig. 4(c) and (d)). It can be observed that there exist twinned a00 martensite variants and amount of martensite variants are distributed in the b matrix, as described in the above corresponding XRD pattern (Fig. 2(d)).

Table 2 Mechanical properties of the studied Ti–24Nb–2Zr, Ti–20Nb–6Zr, Ti–18Nb–8Zr and Ti–16Nb–10Zr alloys. The solution-treated superelastic Ti–22Nb–4Zr is also indicated. Alloys

YS(Mpa)

UTS(Mpa)

ER(%)

E(GPa)

er (%)

scrss (MPa)

Ti26Nb Ti24Nb2Zr Ti22Nb4Zr[12] Ti20Nb6Zr Ti18Nb8Zr Ti16Nb10Zr

390 345 320 365 455 485

470 440 480 470 550 520

19 30 37 39 37 22

65 62 N/A 60 75 70

1.59 1.85 N/A 2.00 1.56 1.65

190 120 160 200 N/A N/A

YS: yield strength; UTS: ultimate tensile strength; ER: elongation at rupture. E: Young’s modulus; er: recovered strain; scrss: critical resolved shear stress.

600 (c)

(d) 500

Stress σb (Mpa)

3.2. Tensile tests Mechanical characterization, including the conventional and cyclic tensile tests, of the solution-treated alloys samples have been performed using tensile tests at room temperature. The measured values determined in this study from the conventional and cyclic tensile tests are reported in Table 2. The conventional tensile curves until fracture of the Ti–(26-z) at% Nb–(z)at% Zr (z ¼0, 2, 6, 8 and 10) alloys are shown in Fig. 5. From the results in our study, it can be seen that an increase of both yield strength for plastic deformation (black arrows) and ultimate tensile strength with the increasing of Zr content via equiatomic substitution Nb element by Zr element. It may be due to the solid-solution hardening by addition of Zr. On the other hand, the stress–strain curves of the four alloys show a large elongation between 39% for Ti–20Nb–6Zr and 22% for

81

(b)

400 (a) Ti22Nb4Zr Ti26Nb

300

200

100

0 0

10

20 Strain (%)

30

40

Fig. 5. Conventional stress–strain curves of Ti–24Nb–2Zr (a), Ti–20Nb–6Zr (b), Ti–18Nb–8Zr (c) and Ti–16Nb–10Zr (d) alloys at room temperature.

Fig. 4. TEM images of the solution-treated Ti–16Nb–10Zr specimen, (a) bright-field image, (b) corresponding diffraction pattern, (c) and (d) dark field image of two a00 variants, respectively.

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600

400

Stress (MPa)

Stress (MPa)

600

E

200

E

200

0

0 0

1

2

3

Strain (%)

4 εr

5

0

1

2

3

Strain (%)

600

4 εr

5

4 εr

5

600

400

Stress (MPa)

Stress (MPa)

400

E 200

0 0

1

2

3

Strain (%)

4

εr

5

E

400

200

0

0

1

2

3 Strain (%)

Fig. 6. Cyclic stress–strain curves of Ti–24Nb–2Zr (a), Ti–20Nb–6Zr (b), Ti–18Nb–8Zr (c) and Ti–16Nb–10Zr (d) alloys at room temperature.

Ti–16Nb–10Zr compared with Ti–26Nb alloy (17%). The Young’s modulus values of the four alloys listed in Table 2 are very low (incipient modulus defined by the slope of the tangent at zero stress) as it is case with many superelastic b-type Ti-based alloys. Especially, the Young’s modulus value of Ti–20Nb–6Zr alloy is only about 60 GPa, which is close to that of bone. This indicates that the Ti–Nb–Zr alloys exhibit low modulus and good ductility. Moreover, the critical stress inducing SIM transformation, scrss, was evaluated (reported in Table 2) and generally increased with the Zr content increasing via equiatomic substitution Nb element by Zr element. The conventional tensile curves (Fig. 5) of the investigated alloys also exhibit two-stage yielding as shown. Concerning the Ti–24Nb–2Zr and Ti–20Nb–6Zr alloys, the microstructure is b and b þ o, the plateau is due to the formation of a00 stress-induced martensite (SIM) during mechanical strain. A similar plateau induced by SIM transformation is also observed in the superelastic b Ti–22Nb–4Zr and Ti–26Nb alloys elaborated by Kim et al. [12]. As a00 martensite is present after quenching in Ti– 16Nb–10Zr (d), the plateau on the stress–strain curve of this alloy is mainly due to the rearrangement of martensite variants. Compared with other three alloys, coexistence of the b phase and a00 martensite are present after quenching in Ti–18Nb–8Zr (c), the plateau on the stress–strain curve of this alloy corresponds to stress-induced transformation of metastable parent phase or the reorientation of martensite present in initial microstructure. The cyclic tensile tests from 0% to 5% strain with strain increments of 0.5% were carried out to evaluate the mechanical properties and reported in Fig. 6. Superelastic recovered strains (er) were measured at 5% strain (10th load–unload cycle) as shown in Fig. 6 and values are reported in Table 2. It can be seen that the superelasticity behavior of the Ti–24Nb–2Zr and Ti–20Nb–6Zr alloys are greater than Ti–18Nb–8Zr and Ti–16Nb–10Zr alloys. The phase constitute of the solution-treated Ti–24Nb–2Zr and Ti–20Nb–6Zr

Fig. 7. TEM images of the as-quenched Ti–20Nb–6Zr specimen after 4% deformation, (a) bright-field image of a single b grain (b) dark field image of stress-induced a00 martensite and corresponding diffraction pattern.

J.Y. Zhang et al. / Materials Science & Engineering A 563 (2013) 78–85

83

Fig. 8. TEM images of the as-quenched Ti–16Nb–10Zr specimen after 4% deformation, (a) bright-field image of the matrix (b þ a00 ) and (b) corresponding diffraction pattern, (c) and (d) dark field image of twinning in the variants a00 and corresponding diffraction pattern.

1000

900

ΔT = Tβ-Ms

2.5

ΔTMax Tβ ΔT

400

300

After the cyclic tests until tensile strains up to 4% of the solution-treated Ti–20Nb–6Zr and Ti–16Nb–10Zr alloys, TEM observations was carried out for the specimens. Fig. 7 shows a typical bright-field TEM image and dark field TEM image with corresponding diffraction pattern of the as-quenched Ti–20Nb–6Zr specimen after 4% deformation. The typical needle-like morphology of a00 martensite phase about 50 nm in width are observed in the b matrix as shown in the bright field presented in Fig. 7(a). While the dark field observations show a00 martensite inside b grains (Fig. 7(b)). The diffraction pattern exhibits typical diffraction spots related to the presence of both a b matrix (higher intensity spot network) and orthorhombic a00 martensite needles (lower intensity spot network) (Fig. 7(b)). This observation confirms that the plateau observed on the tensile curve for this alloy composition is related to the reversible stress-induced a00 martensitic transformation. Low amount of residual SIM a00 also confirms the high reversibility of b/a00 transformation resulting in superior superelasticity, which is in good accordance with the experimental results shown above. Fig. 8 shows bright field TEM image with the corresponding electron diffraction pattern and dark field of variants a00 martensite with the corresponding electron diffraction pattern representing the microstructure of the deformed Ti–16Nb–10Zr alloy. Two deformed variants of a00 martensites described in Fig. 4 were

Ms
ΔT~695K ΔT~630K

~695K

M s>RT

2.0

ΔT~660K 1.5

RT

Recovered strain, εr (%)

3.3. Microstructure after the cyclic deformation

εr

Temperature (K)

alloys samples are fully b phase or b phase and athermal o phase in contrast to the Ti–18Nb–8Zr and Ti–16Nb–10Zr alloys, where the phase constitution are the coexistence of a00 martensite and retain b phases. Therefore, the superelasticity behavior was caused by a stress-induced phase transformation (b-a00 ) during the mechanical test for Ti–24Nb–2Zr and Ti–20Nb–6Zr alloys. For the Ti–18Nb–8Zr and Ti–16Nb–10Zr alloys, the pseudoelastic behaviors mostly resulted from a00 martensite rearrangement as well as the reversible b/a00 martensite transformation from retain b phases during loading.

Measured Ms calculated Ms 0.0

0.1

0.2

0.3 0.4 Zr/Nb ratio

0.5

1.0

0.6

Fig. 9. A schematic plot of Tb, Ms and strain recovery as a function of Zr/Nb ratio for Ti–Nb–Zr alloys.

observed as shown in the bright field image and corresponding diffraction pattern presented in Fig. 8(a) and (b). In detail, the internal twinning involved in the deformed martensitic was also detected in the corresponding SAD pattern (Fig. 8(c) and (d)). The results also indicate that the plateau observed on the tensile curve for this alloy composition is probably related to the reorientation of the martensitic variants occurred by the deformation of twinning.

4. Discussion Equiatomic substitution of Nb element by Zr element in Ti–Nb– Zr alloys results in a linear increase of Tb temperature because of weak influence on Tb for neutral Zr element whereas Nb behaves as a classical b stabilizer in decreasing Tb temperature. In literature,

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the equiatomic substitution kept the Ms temperature of the alloys at almost the same level due to the Ms temperature decreased by 38 K with 1 at% increasing of Zr content and by 40 K with 1 at% increasing of Nb content [12,27]. In fact, the Ms temperature is not only affected by Zr content, but also depends on the Nb concentration. In Ti–Nb–Zr alloys, the Zr influence on measured Ms temperature does not give a simple linear relationship, but exhibits a positive correlation with Nb concentration at certain level [33,34]. Concerning the superelastic property, the Ms temperature is determined at around 240 K, which corresponds to about 26% at Nb content in binary Ti–Nb alloys. After equiatomic substitution of Nb by Zr, the decrease of Nb content results in weakening the Zr ability on suppressing Ms temperature in Ti–Nb–Zr alloys [34]. As observed experimentally in ST specimens, a00 martensites start to appear when Nb content decreased below 20 at%. Therefore, the Zr/Nb ratio is utilized to evaluate the alloying effectiveness on b phase stability by monitoring b transus temperature (Tb), calculated/measured Ms temperature and strain recovery rate. Fig. 9 shows the schematic plot of Tb and Ms as a function of Zr/Nb ratio for Ti–Nb–Zr alloys. The magnitude of calculated Tb and Ms (listed in table 1) increase linearly with the Zr/Nb ratio increase according to the alloy design. However, the experimental results show that equiatomic substitution of Nb element by Zr element has important influence and induces notable changes on the phase constitution in the Ti–Nb–Zr alloys. The phase constitution evolutes from single b phase or b/oath phase to the coexistence of a00 martensite and b phase. The change indicates that the measured Ms temperature raises rapidly when Nb content below the critical value, at which the contribution of Nb becomes inadequate in stabilizing b phase at room temperature. Furthermore, the decreasing Nb concentration weakens the stabilization performance of Zr on keeping the Ms temperature even though the content of Zr is increased. In fact, the experimental magnitude of Ms measured by DSC is appropriate 370 K for Ti–16Nb–10Zr compared with 245 K for Ti–26Nb and 250 K for Ti–20Nb–6Zr. It can be seen that fully b phase or b/oath phase were exhibited when the Zr/Nb ratio is below or equals to 0.3, which corresponds to Ms 250 K. Interesting features should be noticed that the maximum of strain recovery is also obtained around this critical value (shown in Fig. 9). The reason of this phenomenon is considered to be related to the depression of the b phase stability which favors high volume fraction of reversible a00 SIM activation. Since the Ms temperature is appropriately maintained at around 250 K at Zr/Nbo0.4, the change of the b stability is actually related to the increase of Tb temperature, which is resulted from Zr content increasing. Therefore, a coefficient of DT, defined as the difference between Tb and Ms, can be promoted to evaluate the b phase stability on superelastic property of Ti–Nb–Zr alloy with MsoRT. By taking into account of the b/a and b/a00 phase transition temperatures, DT can comprehensively reflect the relationship between b phase stability and the physical background of the metastable alloys. It is consistent that binary Ti–26Nb alloy (Zr/Nb¼ 0) presents normal superelasticity, whereas the Ti–20Nb–6Zr (Zr/Nb¼0.3) exhibits better combined performance (listed in table 2) while the actual DT is up to maximum value. Beyond Zr/Nb  0.42, DT is no longer applicable on superelasticity study due to the measured Ms temperature is close to or higher than room temperature. The corresponding phase constitution and the subsequent main deformation mechanism are altered from reversible a00 SIM transformation to mechanical accommodation of pre-existing a00 martensites. As a result, the theoretical superelastic potential of ternary Ti–Nb–Zr system reaches its upper limit at Nb/Zr ratio around 0.3. It is reasonable to suggest that the high temperature stability of the b phase, indicated by Tb, also plays as an important role as Ms in affecting the superelasticity of Ti–Nb–Zr alloy. Furthermore, the mechanical properties that related to the b

stability, such as elastic modulus and ultimate tensile strength, are sensitive to the DT level in a greater range of metastable b alloys with much lower Ms temperature [38]. Obvious tendency can be noticed that the apparent modulus deceases as Tb increases, which is caused by the creation of a high level of b metastabililty. The improved metastabililty effectively facilitates the occurrence of stress/strain-induced products such as mechanical twinning and phase transformations of a00 martensite and/or o phase. In the design of metastable Ti-alloys, DT is supposed to be indicative based on the current study and some of previous results [10,38]. The addition of neutral and/or interstitial elements, such as Zr, Sn, O and N, even a stabilizer Al can be effective in controlling DT to achieve superior performances on various specific aspects, such as superelasticity, shape memory effect, enhanced ductility and combination of low modulus and high strength.

5. Conclutions In this work, influence of equiatomic Zr/Nb substitution on microstructure evolution, mechanical properties and deformation mechanism of the superelastic Ti–Nb–Zr alloy are investigated. The following results were obtained. (1) The phase constitution change from single b phase or b/oath phase to the coexistence of a00 martensite and b phase with the increasing of Zr content via equiatomic substitution Nb element by Zr element. (2) The mechanism of the superelastic behavior of the studied alloys alters from reversible b/a00 martensitic transformation to rearrangement of pre-existing a00 martensites as increase of Zr content. (3) The measured Ms temperature is maintained at around 250 K with the Zr/Nb ratio changing from 0 to 0.35, then raises rapidly beyond room temperature at Zr/Nb40.4 due to the attenuation of Zr effect in suppressing Ms temperature. (4) The Ti–20Nb–6Zr (Zr/Nb¼0.3) alloy exhibits the best superelasticity and combined performance compared with other three alloys. A coefficient of DT is then proposed to analyze the experimental results by evaluating the b instability of Ti–Nb–Zr alloys, At Zr/Nb  0.3, the maximum value of DT can be obtained, which suggests the superior potential of the theoretical superelasticity of ternary Ti–Nb–Zr system.

Acknowledgments This work was supported by the French National Research Agency (No. ANR 08MAPR 0017).

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