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Characterization of a rapidly annealed Ti50Ni25Cu25 melt-spun ribbon
Journal of Alloys and Compounds 456 (2008) 170–177
Characterization of a rapidly annealed Ti50Ni25Cu25 melt-spun ribbon Yunxiang Tong a , Yong Liu a,∗ , Zeliang Xie b a
School of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore b School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore Received 21 December 2006; received in revised form 11 February 2007; accepted 12 February 2007 Available online 20 February 2007
Abstract In the present paper, crystallization behavior, martensitic transformation and deformation behavior of a Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C under rapid heating were studied. The initially amorphous ribbon was fully crystallized under rapid thermal annealing at 400 ◦ C for 30 s. However, under conventional thermal annealing, crystallization temperature was above 450 ◦ C. The crystallization at low temperature under rapid heating was attributed to the release of the extra energy due to the annihilation of free-volume in amorphous state. With increasing the annealing time, martensitic transformation temperatures increased. The ribbon annealed at 400 ◦ C for 300 s has developed good shape recovery properties. A 2.7% one-way memory strain was obtained the tensile strain of 5.5%. The deformation behavior showed non-flat stress-plateau and high spring-back, which is discussed in terms of grain size and orientation distribution of martensite variants. © 2007 Elsevier B.V. All rights reserved. Keywords: TiNiCu ribbon; Rapid thermal annealing; Crystallization; Martensitic transformation; Deformation behavior; Two-way memory effect
1. Introduction In the past decades, binary NiTi shape memory alloys (SMA) have attracted much attention due to their best combination of various properties. The properties of NiTi alloy can be readily changed by substituting the third element for Ni or Ti. As compared to other alloying elements, the effect of Cu is particular, which leads to much narrower transformation hysteresis, lower martensitic yield strength and less compositional sensitivity of Ms temperature [1]. Melt-spinning technique is an alternative way to prepare almost-ready-to-use NiTi-based alloys and shows several advantages over the conventional techniques [2,3]. The as-spun ribbons are mostly fully or partially amorphous depending on the melt-spinning condition and composition. Annealing is required to enable the as-spun ribbon to transform to crystalline phase in order to have shape memory effect (SME). It has been recently reported that, after proper annealing treatment, significantly different from its bulk counterpart, the Ti50 Ni25 Cu25 melt-ribbon exhibits excellent thermomechanical properties [4–6]. However, the properties of the ribbon are sensitive to annealing
∗
Corresponding author. Tel.: +65 6790 4951; fax: +65 6790 4951. E-mail address: [email protected] (Y. Liu).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.075
temperature and time. Conventional thermal annealing (CTA) typically heats the sample to above 450 ◦ C with a ramping rate lower than 100 ◦ C/min. As compared to the CTA, rapid thermal annealing (RTA) can anneal the ribbon or thin film with a very high heating rate and much shorter dwell time. This favors more precisely control of microstructure and desired properties. The crystallization of the Ti50 Ni25 Cu25 ribbon is governed by polymorphous mechanism and characterized by site-saturation nucleation and isotropic growth [7]. Crystallization kinetics and microstructure of the as-spun Ti50 Ni25 Cu25 ribbon has been reported earlier [2,7,8]. Previous investigations, where the samples were mostly annealed by CTA, show that annealing at a temperature higher than 450 ◦ C is necessary to obtain fully crystallized material [2,8,9]. In the present study, the initially amorphous Ti50 Ni25 Cu25 ribbon was found to be fully crystallized at 400 ◦ C for only 30 s under RTA. The structural evolution as a function of the annealing time was studied. Further, martensitic transformation and deformation behavior of the Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C under RTA was investigated. The results show that no precipitates have formed in the ribbon annealed at 400 ◦ C under RTA, and the rapidly annealed ribbon possesses small transformation hysteresis and good shape memory effect.
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Fig. 1. A typical heating and cooling profile showing the setting temperature and the sample temperature during RTA treatment.
2. Experimental procedures An as-spun Ti50 Ni25 Cu25 (at.%) ribbon with a thickness of about 0.03 mm and width of 23 mm was studied. The as-spun samples were annealed on a Jipelec Jetfirst 100 rapid thermal processor in an argon gas atmosphere. During annealing process, the ribbon was placed on a silicon wafer. The temperature measured by a thermocouple gently contacting the silicon wafer was taken as the sample temperature. The sample was firstly heated up to 200 ◦ C in 10 s and kept for 20 s. Then the sample was heated up to 400 ◦ C at a rate of 50 ◦ C/s and isothermally maintained for different durations from 5 to 300 s. The heating process was controlled precisely to avoid overheating. A typical heating and cooling profile is shown in Fig. 1. A TA instrument 2920 differential scanning calorimeter (DSC) was used to investigate crystallization behavior and martensitic transformation. The transformation characteristics of the ribbon were determined under 5 ◦ C/min cooling and heating rates. Shape memory properties were carried out on an Instron 8800 microforce system equipped with a thermal chamber in which the testing temperature can be changed from −75 to 250 ◦ C at a preset heating/cooling rate. The gauge length of the samples was fixed at 18 mm between the two clamps. Deformation was performed at a strain rate of 2 × 10−5 s−1 . The load cell has a maximum capacity of 50 N with an accuracy of 0.025 N. Pull rods made of quartz bars were used to minimize the effect of thermal expansion. The structure of the Ti50 Ni25 Cu25 ribbon was examined by X-ray diffraction (XRD) using a Philips PW3179 diffractometer with Cu K␣ radiation. All experiments were performed at room temperature. The microstructure was observed on a JEOL 2010 transmission electron microscopy (TEM) operating at 200 kV. Thin foils for TEM observations were prepared by using a Gatan 691 precision ion polishing system (PIPS) at an incidence angle of 4◦ .
Fig. 2. XRD pattern (a) and TEM bright field image (b) of the as-spun Ti50 Ni25 Cu25 ribbon.
ary between amorphous matrix and crystalline phase is smooth and well delineated. The crystallization process of the ribbon was determined using DSC by heating the sample to 550 ◦ C at a rate of 10 ◦ C/min, as shown in Fig. 3. The crystallization temperature was deter-
3. Results and discussion 3.1. As-spun ribbon The microstructure of the as-spun Ti50 Ni25 Cu25 ribbon is characterized to be mainly amorphous with a few crystalline particles embedded in the amorphous matrix. XRD pattern of the as-spun ribbon (Fig. 2(a)) shows mainly a diffuse and low intensity peak of amorphous. Only a few very weak diffraction peaks are observed, that can be indexed as B19 martensite, implying that crystalline particles are very few and they are in martensite phase at room temperature. The existence of some crystalline particles was also confirmed by TEM observation, as shown in Fig. 2(b). Similar to the previous report [2], the bound-
Fig. 3. DSC curve of the as-spun Ti50 Ni25 Cu25 ribbon under 10 ◦ C/min from 40 to 550 ◦ C. The inset is the enlargement of the region around the arrow.
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diffraction peaks from precipitates are visible. These diffraction patterns show that the ribbon directly transformed from amorphous phase into crystalline phase without passing through intermediate phase. The precipitation of the coherent B11–TiCu phase is expected to occur when the Ti50 Ni25 Cu25 alloy is cooled to below 800 ◦ C [1]. For the Ti50 Ni25 Cu25 ribbon, the same precipitates have formed when annealed at 410 ◦ C for a long duration [9]. In order to further confirm the absence of the TiCu precipitates, microstructure observation was carried out. Fig. 5 shows the bright field TEM image and corresponding selected area diffraction pattern (SADP) of the ribbon annealed at 400 ◦ C for 300 s. No TiCu precipitates can be detected, which is also confirmed by the SADP as no streaks due to the thin-plate coherent phase can be observed. Fig. 4. XRD patterns of the as-spun Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 5, 10, 20, 30, 60 and 300 s, respectively.
mined to be 456 ◦ C, close to the value reported previously [2,7]. A weak and diffuse exothermic peak indicated by arrow was also observed before the sharp crystallization peak appears, as shown in the inset of Fig. 3. The origin of this peak is possibly due to the structural relaxation of the amorphous ribbon prior to crystallization. Similar feature was also observed in as-deposited amorphous NiTi thin films [10]. 3.2. Structural evolution Fig. 4 shows the XRD patterns of the ribbon as a function of annealing time at 400 ◦ C under RTA. The structural evolution can be clearly seen. With increasing the annealing time, the wide and diffuse peak representing the amorphous phase fades away and the sharp peaks representing the crystalline phase become distinct. After annealed for 30 s, the ribbon is fully crystallized. All the XRD patterns can only be indexed as B19 martensite ˚ b = 4.297 A, ˚ with the following lattice parameters: a = 2.905 A, ˚ c = 4.527 A, which are close to the reported values [11]. No
3.3. Martensitic transformation The transformation behavior of the Ti50 Ni25 Cu25 ribbon was subsequently studied as a function of the annealing time. Fig. 6 shows the DSC curves for the as-spun sample and those annealed at 400 ◦ C for different time between 5 s and 300 s. Both forward and reverse transformation curves of all samples show multiplestep transformation except for the as-spun sample and the sample annealed for 10 s. The DSC curves of the as-spun ribbon show the weak peaks during cooling and heating, which also indicates that there exist some crystalline particles. This agrees well with the XRD results and TEM observation shown in Fig. 2. When the annealing time was less than 30 s, the weak and wide transformation peaks are found in the DSC curves. On the contrary, the sharp and well-defined DSC peaks are visible after annealed at 400 ◦ C for more than 30 s. The fully crystallized samples show small transformation hysteresis (Af –Ms ), for instance, the transformation hysteresis of the sample annealed for 300 s is 6 ◦ C. The peak transformation temperatures of the samples as a function of the annealing time are plotted in Fig. 7. It is obvious that all the transformation temperatures increase with increasing the annealing time.
Fig. 5. TEM bright field image (a) and the corresponding SAED pattern (b) of the as-spun Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s. The incident electron beam is parallel to the [100]B19 .
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Fig. 6. DSC cooling (a and c) and heating (b and d) curves for the as-spun Ti50 Ni25 Cu25 ribbon and those annealed under RTA at 400 ◦ C for 5, 10, 20, 30, 60 and 300 s, respectively.
3.4. Deformation behavior In the present study, the deformation behavior of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s was further investigated. Fig. 8 plots the stress–strain curves of the sample deformed at room temperature. The testing temperature was about 43 ◦ C below the As temperature, which ensures that the sample was deformed in martensite. The tensile strain was increased subsequently from 1 to 5.5% with an interval of 0.5%. Following each unloading process, the sample was heated to recover the deformed shape and then cooled to room temper-
Fig. 7. Transformation temperatures as a function of the annealing time for the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C.
ature. A same sample was used throughout the test. The present sample fractured before reaching 6% strain. Within this deformation range, the stress–strain curves could be divided into three stages. In the first stage, the stress increases rapidly. In the second stage, the curves are characterized by a non-flat stress-plateau, the stress slowly increases until the strain reaches about 2.5%. In the third stage, the stress increases linearly with increasing the strain. Fig. 9 shows the strain–temperature curve of the sample after deformed to 5.5%. Several characteristic parameters can be defined as follows: εel is the spring-back strain upon unload-
Fig. 8. Stress–strain curves of the Ti50 Ni25 Cu25 ribbon deformed at room temperature. The ribbon was annealed under RTA at 400 ◦ C for 300 s.
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Fig. 9. Temperature–strain curves of the ribbon annealed under RTA at 400 ◦ C for 300 s after deformed to 5.5% at room temperature.
Fig. 11. Effect of the tensile strain on the two-way recovery strain (εtw ). The ribbon was annealed under RTA at 400 ◦ C for 300 s.
ing, εre is the one-way shape recovery strain upon heating, εir is the irreversible strain after heating and εtw is the two-way memory strain after cooling to martensite. It is seen that the sample shrank slowly at the beginning of the heating. When the temperature was higher than As , the shape recovery proceeded rapidly. An obvious two-way memory effect (TWME) was induced upon cooling. Fig. 10 shows the effect of the tensile strain on εel , εre and εir . The contribution of thermal expansion has been removed. With increasing the tensile strain, εel , εre and εir increase. When the tensile strain was less than 2%, it can completely recover its original shape upon heating. Beyond 2% deformation, an irreversible strain presents because the dislocations are introduced. εel increases slowly at the initial stage and then progresses rapidly, which is different from the trend of εre . The dashed line plots the linear elasticity evolution with the tensile strain. The deviation of εel from linearity is due to the dislocation formation. A maximum one-way recovery strain is about 2.7% with an irreversible strain of 0.54% after 5.5% deformation. Fig. 11 shows the effect of the tensile strain on εtw . It is seen that the present sample does not exhibit any noticeable TWME when it was tensioned to less than 2.5% strain (end of
the stress-plateau). This is different from the TWME developed in the equiatomic NiTi alloy through martensite deformation, in which a 1.5% εtw was observed at 6.2% tensile strain corresponding to the end of the stress-plateau [12]. When the tensile deformation exceeds the limit of the stress-plateau, εtw increases as expected. A maximum εtw of 0.56% was obtained at 5.5% tensile-strain, which is lower than that of equiatomic NiTi alloy [12] and Ti50 Ni40 Cu10 alloy [13] and higher than solution-treated Ni49 Ti35 Hf15 alloy [14]. Fig. 12 shows the transformation behavior of the Ti50 Ni25 Cu25 ribbon after deformed to 5.5% strain at room temperature. The transformation interval upon first heating is much smaller as compared to that during the second heating. The reverse transformation start temperature (As2 ) is 60.4 ◦ C upon the second heating, about 4 ◦ C lower than that in the first heating. It implies that the martensite is slightly stabilized by the deformation. The difference between the As temperatures is much lower than that for equiatomic NiTi alloy (about 20 ◦ C) [12] and Ni49 Ti36 Hf15 alloy (about 13 ◦ C) [14] with the same tensile strain. This agrees with the deformation behavior of less dislocation process. The As temperature upon first heating as a function of the tensile strain is shown in Fig. 13. The transforma-
Fig. 10. Effect of the tensile strain on the elastic spring-back strain (εel ), oneway recovery strain (εre ) and irreversible strain (εir ). The ribbon was annealed under RTA at 400 ◦ C for 300 s.
Fig. 12. DSC curve of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s after deformed to 5.5% at room temperature.
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Fig. 13. Effect of the tensile strain on the reverse transformation temperatures of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s.
tion temperatures were determined from the strain–temperature curves shown in Fig. 9 by using cross-tangential line method. The As temperature continuously increases with increasing the tensile strain when the deformation is beyond the limit of the stress-plateau. Comparing the results shown in Figs. 6 and 12, it is found that the transformation hysteresis of the deformed Ti50 Ni25 Cu25 ribbon is almost the same as the undeformed sample, however, the transformation interval of the former is larger than that of the latter. 3.5. Discussion 3.5.1. Crystallization behavior The full crystallization of the Ti50 Ni25 Cu25 ribbon was normally achieved by CTA at a temperature higher than 450 ◦ C [2]. It remains the initial amorphous structure after annealed at 400 ◦ C for 30 min under CTA [15]. However, in the present study, the initially amorphous ribbon was fully crystallized by RTA at 400 ◦ C for 30 s, as evidenced by XRD tests. Although R¨osner et al. once reported that annealing at 410 ◦ C under CTA can fully crystallize the ribbon, it has to take at least 24 h [9]. The precipitation of TiCu phase unavoidably occurs during such a long annealing time, which deteriorates the mechanical and thermomechanical properties of the Ti50 Ni25 Cu25 ribbon [4,15,16]. It must be emphasized that no precipitates were found in the present low temperature crystallized sample. This can be attributed to the combination of lower annealing temperature and shorter annealing duration. It is important to understand the crystallization path that begins with amorphous phase and ends with B2 parent phase. In the as-spun ribbon, the atoms are randomly arranged with excess energy stored in the metastable structure. Structural relaxation usually occurs during annealing and changes the metastable state to the state with lower energy [17]. In the present case, the transformation path is likely the amorphous-to-structural relaxation-to-crystalline phase. The as-spun amorphous ribbon contains a large number of defects that are mainly free volume [18] and short-range-ordered clusters. Structural relaxation is usually associated with the annihilation of the free volume characterized by the wide and diffuse
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Fig. 14. DSC curve (10 ◦ C/min) of the as-spun Ti50 Ni25 Cu25 ribbon preannealed under CTA at 400 ◦ C for 15 min. The inset highlights the details around the crystallization temperature of the curve.
exothermal peak in DSC curve of Fig. 3. After crystallization, the free volume can be totally annealed out. In order to investigate the effect of the structure relaxation on crystallization, a sample was heated up to 400 ◦ C with a heating rate of 10 ◦ C/min and then kept isothermally for 15 min. XRD results show that the sample was still amorphous, same as the as-spun ribbon. This pre-annealed sample was then tested in DSC, and the results are shown in Fig. 14. Comparing the results shown in Figs. 3 and 14, no weak and diffuse peak was observed prior to crystallization, indicating no obvious structure relaxation in the second heating. This pre-annealed sample shows identical crystallization temperature with the as-spun ribbon, suggesting that the structural relaxation does not have obvious effect on the crystallization under CTA. The above pre-annealed sample was also heated under RTA from room temperature to 400 ◦ C and remained for 30 s. The XRD result shown in Fig. 15(a) reveals that the sample was partially crystallized, different from the result of the asspun ribbon directly annealed under RTA shown in Fig. 15(b). This suggests that the crystallization of the ribbon became more
Fig. 15. Comparison of the XRD patterns of the Ti50 Ni25 Cu25 samples. (a) The as-spun ribbon was annealed under CTA at 400 ◦ C for 15 min followed by annealed under RTA at 400 ◦ C for 30 s. (b) The as-spun ribbon was directly annealed under RTA at 400 ◦ C for 30 s.
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difficult after structure relaxation. During RTA, the structural relaxation was shifted to higher temperature because of the rapid heating. The excess energy due to a higher internal stress field associated with free-volume can provide an additional driving force for crystallization. It has been found that some short-range-ordered clusters exist in the as-spun Ti50 Ni25 Cu25 ribbon [19]. This may result from the fact that Ti–Ni pair has a much larger mixing enthalpy than Ti–Cu and Ni–Cu pairs in the present alloy (Ti–Ni: −35 kJ/mol, Ti–Cu:−9 kJ/mol, and Ni–Cu: +4 kJ/mol) [20]. Thus, Ti should have strong interaction with Ni, which possibly leads to the formation of short-range-ordered clusters of Ti–Ni pair in the liquid state. During melt-spinning, the metastable atomic configuration is frozen in the amorphous matrix. During RTA, the short-range-ordered clusters can act as the nucleation sites. In summary, the low temperature crystallization under RTA is naturally related to the effect of the heating rate, as compared to the crystallization under CTA. The structural relaxation may assist the crystallization by providing an extra energy. The low temperature crystallization under RTA may reduce the processing cost as compared to the crystallization under CTA. Such processing is also beneficial to obtain the desired microstructure, such as grain size and size of the precipitates, thus controlling the properties. 3.5.2. Martensitic transformation From the DSC results shown in Figs. 6 and 7, martensitic transformation behavior is strongly dependent on the annealing conditions. Several microstructural factors should be considered notably, amorphous matrix, grain size and precipitates. The effect of the precipitates can be firstly ruled out since no precipitates were found in the annealed ribbon. The effect of amorphous-crystalline interface on the martensitic transformation is the same as that of a smaller grain size [21,22]. Both hinder the transformation by increasing the non-chemical energy terms. When annealed for less than 30s, the samples were only partially crystallized. Therefore, the complex transformation is related to the combined effect arising from the amorphous-crystalline interface and the difference in grain size. After annealed for more than 30 s, the samples were fully crystallized. The newly formed grains usually have a much smaller size than the preexisted grains formed during rapid solidification [2,3,23]. The two-step martensitic transformation could be ascribed to this grain size difference. With increasing the annealing time, the grain size increases, which causes the transformation shifted to higher temperature side. Similar effect of grain size on the martensitic transformation has been reported in other NiTi-based alloys [22]. 3.5.3. Deformation behavior The deformation of the present Ti50 Ni25 Cu25 sample proceeds in three different stages. The first stage corresponds to the elastic deformation of self-accommodated martensite variants. In the second stage, the deformation is mainly related to the reorientation/detwinning of the martensite variants. The martensite variants are easily reoriented by the external force, as evidenced by the relatively low stress to induce the martensite reorienta-
tion (about 60 MPa). In this stage, the sample exhibits the largest recovery ratio. In the third stage, the recovery strain continuously increases as shown in Fig. 10, indicating that a further detwinning of the martensite variants occurs. Meanwhile, the irreversible strain also increases with increasing the strain, indicating the generation of dislocations. The above features are quite similar to the equiatomic NiTi alloy [12,24]. Several factors can be related to the deformation behavior of the rapidly annealed Ti50 Ni25 Cu25 ribbon, including grain size, texture and martensite structure. The non-flat stress-plateau implies that the detwinning process of the Ti50 Ni25 Cu25 ribbon requires continuous increase in stress to provide the driving force. This may be related to grain size effect. Most of the grains in the present Ti50 Ni25 Cu25 ribbon have sizes less than 1 m, much smaller than that of conventionally prepared alloys. It is known that the detwinning process proceeds unevenly within the plateau region [25]. The localized internal stress is unable to trigger the detwinning of the neighboring martensite twins of less favorably oriented because of the constraint from the grain boundaries. This is also supported by the deformation behavior in NiTi thin film with ultrafine grains [26]. The twinning modes in B19 martensite of Ti50 Ni25 Cu25 ribbon are (1 1 1) type I twin and (0 1 1) compound twin. For (1 1 1) type I twins, the twinning plane is (1 1 1) with shear direction [1.1397 0.1397 1]. For (0 1 1) compound twins, the twinning ¯ [2,15]. It is found plane is (0 1 1) plane with shear direction [011] that the “just crystallized” ribbon has a pure [211] fiber texture with its fiber axis perpendicular to the ribbon surface [15,27]. The (1 1 1) twinning plane has an angle of about 5◦ to the ribbon surface, whereas for the (0 1 1) compound twinning plane the angle is about 47◦ . Both the shear directions randomly distribute within the twinning planes. The deformation process in the low temperature crystallized Ti50 Ni25 Cu25 ribbon involves mainly the detwinning of the (1 1 1) type I twin which is isotropic due to the fiber texture. Such isotropic detwinning results in smaller plateau strain and earlier dislocation process. This further leads to lower shape memory strain and high spring back. The TWME developed in Ti50 Ni25 Cu25 ribbon is related to the specially arranged dislocations produced during martensite deformation. The dislocations would create the directional internal stress field and causes preferential growth of martensite variants [12]. This low TWME should be related to the fiber texture that leads to formation of randomly oriented of martensite variants. In addition, the sample with a smaller grain size usually has a higher yield stress and suppresses the generation of dislocations. This further results in less preferentially oriented martensite variants upon cooling, thus, smaller two-way memory strain. 4. Conclusions 1. The initially amorphous ribbon is fully crystallized under RTA at 400 ◦ C for 30 s. However, under CTA the ribbon remains amorphous even after annealed at 400 ◦ C for 30 min. The crystallization at low temperature under rapid heating is attributed to the release of the extra energy stored in the free-volume of amorphous state.
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2. No precipitates of B11–TiCu phase are found in the Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C for up to 300 s under RTA. 3. With increasing the annealing time, the transformation temperatures increase. The dependence of transformation behavior on the annealing time could be related to the constraint from the amorphous–crystalline interface and the grain boundaries. 4. For the sample annealed at 400 ◦ C for 300 s, a one-way recovery strain of 2.7% is obtained with the tensile strain of 5.5%. The stress–strain curves show non-flat stress-plateau and high spring back. A TWME is induced by martensite deformation. Both the deformation behavior and TWME are likely related to the isotropic detwinning as a result of formation of randomly oriented martensite variants. The small grain size also contributes to the large spring back and low TWME. References [1] W.J. Moberly, K.N. Melton, in: T.W. Duerig, K.N. Melton, D. St¨okel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann Ltd., London, UK, 1990, p. 46. [2] Z.L. Xie, J. Van Humbeeck, Y. Liu, L. Dalaey, Scripta Mater. 37 (1997) 363. [3] Yeon-Wook Kim, Tae-Hyun Nam, Scripta Mater. 51 (2004) 653. [4] Y. Liu, Z.L. Xie, Y.X. Tong, C.W. Lim, J. Alloys Compd. 416 (2006) 188. [5] Y. Liu, Mater. Sci. Eng. A 354 (2003) 286.
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Characterization of a rapidly annealed Ti50Ni25Cu25 melt-spun ribbon Yunxiang Tong a , Yong Liu a,∗ , Zeliang Xie b a
School of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore b School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore Received 21 December 2006; received in revised form 11 February 2007; accepted 12 February 2007 Available online 20 February 2007
Abstract In the present paper, crystallization behavior, martensitic transformation and deformation behavior of a Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C under rapid heating were studied. The initially amorphous ribbon was fully crystallized under rapid thermal annealing at 400 ◦ C for 30 s. However, under conventional thermal annealing, crystallization temperature was above 450 ◦ C. The crystallization at low temperature under rapid heating was attributed to the release of the extra energy due to the annihilation of free-volume in amorphous state. With increasing the annealing time, martensitic transformation temperatures increased. The ribbon annealed at 400 ◦ C for 300 s has developed good shape recovery properties. A 2.7% one-way memory strain was obtained the tensile strain of 5.5%. The deformation behavior showed non-flat stress-plateau and high spring-back, which is discussed in terms of grain size and orientation distribution of martensite variants. © 2007 Elsevier B.V. All rights reserved. Keywords: TiNiCu ribbon; Rapid thermal annealing; Crystallization; Martensitic transformation; Deformation behavior; Two-way memory effect
1. Introduction In the past decades, binary NiTi shape memory alloys (SMA) have attracted much attention due to their best combination of various properties. The properties of NiTi alloy can be readily changed by substituting the third element for Ni or Ti. As compared to other alloying elements, the effect of Cu is particular, which leads to much narrower transformation hysteresis, lower martensitic yield strength and less compositional sensitivity of Ms temperature [1]. Melt-spinning technique is an alternative way to prepare almost-ready-to-use NiTi-based alloys and shows several advantages over the conventional techniques [2,3]. The as-spun ribbons are mostly fully or partially amorphous depending on the melt-spinning condition and composition. Annealing is required to enable the as-spun ribbon to transform to crystalline phase in order to have shape memory effect (SME). It has been recently reported that, after proper annealing treatment, significantly different from its bulk counterpart, the Ti50 Ni25 Cu25 melt-ribbon exhibits excellent thermomechanical properties [4–6]. However, the properties of the ribbon are sensitive to annealing
∗
Corresponding author. Tel.: +65 6790 4951; fax: +65 6790 4951. E-mail address: [email protected] (Y. Liu).
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temperature and time. Conventional thermal annealing (CTA) typically heats the sample to above 450 ◦ C with a ramping rate lower than 100 ◦ C/min. As compared to the CTA, rapid thermal annealing (RTA) can anneal the ribbon or thin film with a very high heating rate and much shorter dwell time. This favors more precisely control of microstructure and desired properties. The crystallization of the Ti50 Ni25 Cu25 ribbon is governed by polymorphous mechanism and characterized by site-saturation nucleation and isotropic growth [7]. Crystallization kinetics and microstructure of the as-spun Ti50 Ni25 Cu25 ribbon has been reported earlier [2,7,8]. Previous investigations, where the samples were mostly annealed by CTA, show that annealing at a temperature higher than 450 ◦ C is necessary to obtain fully crystallized material [2,8,9]. In the present study, the initially amorphous Ti50 Ni25 Cu25 ribbon was found to be fully crystallized at 400 ◦ C for only 30 s under RTA. The structural evolution as a function of the annealing time was studied. Further, martensitic transformation and deformation behavior of the Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C under RTA was investigated. The results show that no precipitates have formed in the ribbon annealed at 400 ◦ C under RTA, and the rapidly annealed ribbon possesses small transformation hysteresis and good shape memory effect.
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Fig. 1. A typical heating and cooling profile showing the setting temperature and the sample temperature during RTA treatment.
2. Experimental procedures An as-spun Ti50 Ni25 Cu25 (at.%) ribbon with a thickness of about 0.03 mm and width of 23 mm was studied. The as-spun samples were annealed on a Jipelec Jetfirst 100 rapid thermal processor in an argon gas atmosphere. During annealing process, the ribbon was placed on a silicon wafer. The temperature measured by a thermocouple gently contacting the silicon wafer was taken as the sample temperature. The sample was firstly heated up to 200 ◦ C in 10 s and kept for 20 s. Then the sample was heated up to 400 ◦ C at a rate of 50 ◦ C/s and isothermally maintained for different durations from 5 to 300 s. The heating process was controlled precisely to avoid overheating. A typical heating and cooling profile is shown in Fig. 1. A TA instrument 2920 differential scanning calorimeter (DSC) was used to investigate crystallization behavior and martensitic transformation. The transformation characteristics of the ribbon were determined under 5 ◦ C/min cooling and heating rates. Shape memory properties were carried out on an Instron 8800 microforce system equipped with a thermal chamber in which the testing temperature can be changed from −75 to 250 ◦ C at a preset heating/cooling rate. The gauge length of the samples was fixed at 18 mm between the two clamps. Deformation was performed at a strain rate of 2 × 10−5 s−1 . The load cell has a maximum capacity of 50 N with an accuracy of 0.025 N. Pull rods made of quartz bars were used to minimize the effect of thermal expansion. The structure of the Ti50 Ni25 Cu25 ribbon was examined by X-ray diffraction (XRD) using a Philips PW3179 diffractometer with Cu K␣ radiation. All experiments were performed at room temperature. The microstructure was observed on a JEOL 2010 transmission electron microscopy (TEM) operating at 200 kV. Thin foils for TEM observations were prepared by using a Gatan 691 precision ion polishing system (PIPS) at an incidence angle of 4◦ .
Fig. 2. XRD pattern (a) and TEM bright field image (b) of the as-spun Ti50 Ni25 Cu25 ribbon.
ary between amorphous matrix and crystalline phase is smooth and well delineated. The crystallization process of the ribbon was determined using DSC by heating the sample to 550 ◦ C at a rate of 10 ◦ C/min, as shown in Fig. 3. The crystallization temperature was deter-
3. Results and discussion 3.1. As-spun ribbon The microstructure of the as-spun Ti50 Ni25 Cu25 ribbon is characterized to be mainly amorphous with a few crystalline particles embedded in the amorphous matrix. XRD pattern of the as-spun ribbon (Fig. 2(a)) shows mainly a diffuse and low intensity peak of amorphous. Only a few very weak diffraction peaks are observed, that can be indexed as B19 martensite, implying that crystalline particles are very few and they are in martensite phase at room temperature. The existence of some crystalline particles was also confirmed by TEM observation, as shown in Fig. 2(b). Similar to the previous report [2], the bound-
Fig. 3. DSC curve of the as-spun Ti50 Ni25 Cu25 ribbon under 10 ◦ C/min from 40 to 550 ◦ C. The inset is the enlargement of the region around the arrow.
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diffraction peaks from precipitates are visible. These diffraction patterns show that the ribbon directly transformed from amorphous phase into crystalline phase without passing through intermediate phase. The precipitation of the coherent B11–TiCu phase is expected to occur when the Ti50 Ni25 Cu25 alloy is cooled to below 800 ◦ C [1]. For the Ti50 Ni25 Cu25 ribbon, the same precipitates have formed when annealed at 410 ◦ C for a long duration [9]. In order to further confirm the absence of the TiCu precipitates, microstructure observation was carried out. Fig. 5 shows the bright field TEM image and corresponding selected area diffraction pattern (SADP) of the ribbon annealed at 400 ◦ C for 300 s. No TiCu precipitates can be detected, which is also confirmed by the SADP as no streaks due to the thin-plate coherent phase can be observed. Fig. 4. XRD patterns of the as-spun Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 5, 10, 20, 30, 60 and 300 s, respectively.
mined to be 456 ◦ C, close to the value reported previously [2,7]. A weak and diffuse exothermic peak indicated by arrow was also observed before the sharp crystallization peak appears, as shown in the inset of Fig. 3. The origin of this peak is possibly due to the structural relaxation of the amorphous ribbon prior to crystallization. Similar feature was also observed in as-deposited amorphous NiTi thin films [10]. 3.2. Structural evolution Fig. 4 shows the XRD patterns of the ribbon as a function of annealing time at 400 ◦ C under RTA. The structural evolution can be clearly seen. With increasing the annealing time, the wide and diffuse peak representing the amorphous phase fades away and the sharp peaks representing the crystalline phase become distinct. After annealed for 30 s, the ribbon is fully crystallized. All the XRD patterns can only be indexed as B19 martensite ˚ b = 4.297 A, ˚ with the following lattice parameters: a = 2.905 A, ˚ c = 4.527 A, which are close to the reported values [11]. No
3.3. Martensitic transformation The transformation behavior of the Ti50 Ni25 Cu25 ribbon was subsequently studied as a function of the annealing time. Fig. 6 shows the DSC curves for the as-spun sample and those annealed at 400 ◦ C for different time between 5 s and 300 s. Both forward and reverse transformation curves of all samples show multiplestep transformation except for the as-spun sample and the sample annealed for 10 s. The DSC curves of the as-spun ribbon show the weak peaks during cooling and heating, which also indicates that there exist some crystalline particles. This agrees well with the XRD results and TEM observation shown in Fig. 2. When the annealing time was less than 30 s, the weak and wide transformation peaks are found in the DSC curves. On the contrary, the sharp and well-defined DSC peaks are visible after annealed at 400 ◦ C for more than 30 s. The fully crystallized samples show small transformation hysteresis (Af –Ms ), for instance, the transformation hysteresis of the sample annealed for 300 s is 6 ◦ C. The peak transformation temperatures of the samples as a function of the annealing time are plotted in Fig. 7. It is obvious that all the transformation temperatures increase with increasing the annealing time.
Fig. 5. TEM bright field image (a) and the corresponding SAED pattern (b) of the as-spun Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s. The incident electron beam is parallel to the [100]B19 .
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Fig. 6. DSC cooling (a and c) and heating (b and d) curves for the as-spun Ti50 Ni25 Cu25 ribbon and those annealed under RTA at 400 ◦ C for 5, 10, 20, 30, 60 and 300 s, respectively.
3.4. Deformation behavior In the present study, the deformation behavior of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s was further investigated. Fig. 8 plots the stress–strain curves of the sample deformed at room temperature. The testing temperature was about 43 ◦ C below the As temperature, which ensures that the sample was deformed in martensite. The tensile strain was increased subsequently from 1 to 5.5% with an interval of 0.5%. Following each unloading process, the sample was heated to recover the deformed shape and then cooled to room temper-
Fig. 7. Transformation temperatures as a function of the annealing time for the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C.
ature. A same sample was used throughout the test. The present sample fractured before reaching 6% strain. Within this deformation range, the stress–strain curves could be divided into three stages. In the first stage, the stress increases rapidly. In the second stage, the curves are characterized by a non-flat stress-plateau, the stress slowly increases until the strain reaches about 2.5%. In the third stage, the stress increases linearly with increasing the strain. Fig. 9 shows the strain–temperature curve of the sample after deformed to 5.5%. Several characteristic parameters can be defined as follows: εel is the spring-back strain upon unload-
Fig. 8. Stress–strain curves of the Ti50 Ni25 Cu25 ribbon deformed at room temperature. The ribbon was annealed under RTA at 400 ◦ C for 300 s.
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Fig. 9. Temperature–strain curves of the ribbon annealed under RTA at 400 ◦ C for 300 s after deformed to 5.5% at room temperature.
Fig. 11. Effect of the tensile strain on the two-way recovery strain (εtw ). The ribbon was annealed under RTA at 400 ◦ C for 300 s.
ing, εre is the one-way shape recovery strain upon heating, εir is the irreversible strain after heating and εtw is the two-way memory strain after cooling to martensite. It is seen that the sample shrank slowly at the beginning of the heating. When the temperature was higher than As , the shape recovery proceeded rapidly. An obvious two-way memory effect (TWME) was induced upon cooling. Fig. 10 shows the effect of the tensile strain on εel , εre and εir . The contribution of thermal expansion has been removed. With increasing the tensile strain, εel , εre and εir increase. When the tensile strain was less than 2%, it can completely recover its original shape upon heating. Beyond 2% deformation, an irreversible strain presents because the dislocations are introduced. εel increases slowly at the initial stage and then progresses rapidly, which is different from the trend of εre . The dashed line plots the linear elasticity evolution with the tensile strain. The deviation of εel from linearity is due to the dislocation formation. A maximum one-way recovery strain is about 2.7% with an irreversible strain of 0.54% after 5.5% deformation. Fig. 11 shows the effect of the tensile strain on εtw . It is seen that the present sample does not exhibit any noticeable TWME when it was tensioned to less than 2.5% strain (end of
the stress-plateau). This is different from the TWME developed in the equiatomic NiTi alloy through martensite deformation, in which a 1.5% εtw was observed at 6.2% tensile strain corresponding to the end of the stress-plateau [12]. When the tensile deformation exceeds the limit of the stress-plateau, εtw increases as expected. A maximum εtw of 0.56% was obtained at 5.5% tensile-strain, which is lower than that of equiatomic NiTi alloy [12] and Ti50 Ni40 Cu10 alloy [13] and higher than solution-treated Ni49 Ti35 Hf15 alloy [14]. Fig. 12 shows the transformation behavior of the Ti50 Ni25 Cu25 ribbon after deformed to 5.5% strain at room temperature. The transformation interval upon first heating is much smaller as compared to that during the second heating. The reverse transformation start temperature (As2 ) is 60.4 ◦ C upon the second heating, about 4 ◦ C lower than that in the first heating. It implies that the martensite is slightly stabilized by the deformation. The difference between the As temperatures is much lower than that for equiatomic NiTi alloy (about 20 ◦ C) [12] and Ni49 Ti36 Hf15 alloy (about 13 ◦ C) [14] with the same tensile strain. This agrees with the deformation behavior of less dislocation process. The As temperature upon first heating as a function of the tensile strain is shown in Fig. 13. The transforma-
Fig. 10. Effect of the tensile strain on the elastic spring-back strain (εel ), oneway recovery strain (εre ) and irreversible strain (εir ). The ribbon was annealed under RTA at 400 ◦ C for 300 s.
Fig. 12. DSC curve of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s after deformed to 5.5% at room temperature.
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Fig. 13. Effect of the tensile strain on the reverse transformation temperatures of the Ti50 Ni25 Cu25 ribbon annealed under RTA at 400 ◦ C for 300 s.
tion temperatures were determined from the strain–temperature curves shown in Fig. 9 by using cross-tangential line method. The As temperature continuously increases with increasing the tensile strain when the deformation is beyond the limit of the stress-plateau. Comparing the results shown in Figs. 6 and 12, it is found that the transformation hysteresis of the deformed Ti50 Ni25 Cu25 ribbon is almost the same as the undeformed sample, however, the transformation interval of the former is larger than that of the latter. 3.5. Discussion 3.5.1. Crystallization behavior The full crystallization of the Ti50 Ni25 Cu25 ribbon was normally achieved by CTA at a temperature higher than 450 ◦ C [2]. It remains the initial amorphous structure after annealed at 400 ◦ C for 30 min under CTA [15]. However, in the present study, the initially amorphous ribbon was fully crystallized by RTA at 400 ◦ C for 30 s, as evidenced by XRD tests. Although R¨osner et al. once reported that annealing at 410 ◦ C under CTA can fully crystallize the ribbon, it has to take at least 24 h [9]. The precipitation of TiCu phase unavoidably occurs during such a long annealing time, which deteriorates the mechanical and thermomechanical properties of the Ti50 Ni25 Cu25 ribbon [4,15,16]. It must be emphasized that no precipitates were found in the present low temperature crystallized sample. This can be attributed to the combination of lower annealing temperature and shorter annealing duration. It is important to understand the crystallization path that begins with amorphous phase and ends with B2 parent phase. In the as-spun ribbon, the atoms are randomly arranged with excess energy stored in the metastable structure. Structural relaxation usually occurs during annealing and changes the metastable state to the state with lower energy [17]. In the present case, the transformation path is likely the amorphous-to-structural relaxation-to-crystalline phase. The as-spun amorphous ribbon contains a large number of defects that are mainly free volume [18] and short-range-ordered clusters. Structural relaxation is usually associated with the annihilation of the free volume characterized by the wide and diffuse
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Fig. 14. DSC curve (10 ◦ C/min) of the as-spun Ti50 Ni25 Cu25 ribbon preannealed under CTA at 400 ◦ C for 15 min. The inset highlights the details around the crystallization temperature of the curve.
exothermal peak in DSC curve of Fig. 3. After crystallization, the free volume can be totally annealed out. In order to investigate the effect of the structure relaxation on crystallization, a sample was heated up to 400 ◦ C with a heating rate of 10 ◦ C/min and then kept isothermally for 15 min. XRD results show that the sample was still amorphous, same as the as-spun ribbon. This pre-annealed sample was then tested in DSC, and the results are shown in Fig. 14. Comparing the results shown in Figs. 3 and 14, no weak and diffuse peak was observed prior to crystallization, indicating no obvious structure relaxation in the second heating. This pre-annealed sample shows identical crystallization temperature with the as-spun ribbon, suggesting that the structural relaxation does not have obvious effect on the crystallization under CTA. The above pre-annealed sample was also heated under RTA from room temperature to 400 ◦ C and remained for 30 s. The XRD result shown in Fig. 15(a) reveals that the sample was partially crystallized, different from the result of the asspun ribbon directly annealed under RTA shown in Fig. 15(b). This suggests that the crystallization of the ribbon became more
Fig. 15. Comparison of the XRD patterns of the Ti50 Ni25 Cu25 samples. (a) The as-spun ribbon was annealed under CTA at 400 ◦ C for 15 min followed by annealed under RTA at 400 ◦ C for 30 s. (b) The as-spun ribbon was directly annealed under RTA at 400 ◦ C for 30 s.
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difficult after structure relaxation. During RTA, the structural relaxation was shifted to higher temperature because of the rapid heating. The excess energy due to a higher internal stress field associated with free-volume can provide an additional driving force for crystallization. It has been found that some short-range-ordered clusters exist in the as-spun Ti50 Ni25 Cu25 ribbon [19]. This may result from the fact that Ti–Ni pair has a much larger mixing enthalpy than Ti–Cu and Ni–Cu pairs in the present alloy (Ti–Ni: −35 kJ/mol, Ti–Cu:−9 kJ/mol, and Ni–Cu: +4 kJ/mol) [20]. Thus, Ti should have strong interaction with Ni, which possibly leads to the formation of short-range-ordered clusters of Ti–Ni pair in the liquid state. During melt-spinning, the metastable atomic configuration is frozen in the amorphous matrix. During RTA, the short-range-ordered clusters can act as the nucleation sites. In summary, the low temperature crystallization under RTA is naturally related to the effect of the heating rate, as compared to the crystallization under CTA. The structural relaxation may assist the crystallization by providing an extra energy. The low temperature crystallization under RTA may reduce the processing cost as compared to the crystallization under CTA. Such processing is also beneficial to obtain the desired microstructure, such as grain size and size of the precipitates, thus controlling the properties. 3.5.2. Martensitic transformation From the DSC results shown in Figs. 6 and 7, martensitic transformation behavior is strongly dependent on the annealing conditions. Several microstructural factors should be considered notably, amorphous matrix, grain size and precipitates. The effect of the precipitates can be firstly ruled out since no precipitates were found in the annealed ribbon. The effect of amorphous-crystalline interface on the martensitic transformation is the same as that of a smaller grain size [21,22]. Both hinder the transformation by increasing the non-chemical energy terms. When annealed for less than 30s, the samples were only partially crystallized. Therefore, the complex transformation is related to the combined effect arising from the amorphous-crystalline interface and the difference in grain size. After annealed for more than 30 s, the samples were fully crystallized. The newly formed grains usually have a much smaller size than the preexisted grains formed during rapid solidification [2,3,23]. The two-step martensitic transformation could be ascribed to this grain size difference. With increasing the annealing time, the grain size increases, which causes the transformation shifted to higher temperature side. Similar effect of grain size on the martensitic transformation has been reported in other NiTi-based alloys [22]. 3.5.3. Deformation behavior The deformation of the present Ti50 Ni25 Cu25 sample proceeds in three different stages. The first stage corresponds to the elastic deformation of self-accommodated martensite variants. In the second stage, the deformation is mainly related to the reorientation/detwinning of the martensite variants. The martensite variants are easily reoriented by the external force, as evidenced by the relatively low stress to induce the martensite reorienta-
tion (about 60 MPa). In this stage, the sample exhibits the largest recovery ratio. In the third stage, the recovery strain continuously increases as shown in Fig. 10, indicating that a further detwinning of the martensite variants occurs. Meanwhile, the irreversible strain also increases with increasing the strain, indicating the generation of dislocations. The above features are quite similar to the equiatomic NiTi alloy [12,24]. Several factors can be related to the deformation behavior of the rapidly annealed Ti50 Ni25 Cu25 ribbon, including grain size, texture and martensite structure. The non-flat stress-plateau implies that the detwinning process of the Ti50 Ni25 Cu25 ribbon requires continuous increase in stress to provide the driving force. This may be related to grain size effect. Most of the grains in the present Ti50 Ni25 Cu25 ribbon have sizes less than 1 m, much smaller than that of conventionally prepared alloys. It is known that the detwinning process proceeds unevenly within the plateau region [25]. The localized internal stress is unable to trigger the detwinning of the neighboring martensite twins of less favorably oriented because of the constraint from the grain boundaries. This is also supported by the deformation behavior in NiTi thin film with ultrafine grains [26]. The twinning modes in B19 martensite of Ti50 Ni25 Cu25 ribbon are (1 1 1) type I twin and (0 1 1) compound twin. For (1 1 1) type I twins, the twinning plane is (1 1 1) with shear direction [1.1397 0.1397 1]. For (0 1 1) compound twins, the twinning ¯ [2,15]. It is found plane is (0 1 1) plane with shear direction [011] that the “just crystallized” ribbon has a pure [211] fiber texture with its fiber axis perpendicular to the ribbon surface [15,27]. The (1 1 1) twinning plane has an angle of about 5◦ to the ribbon surface, whereas for the (0 1 1) compound twinning plane the angle is about 47◦ . Both the shear directions randomly distribute within the twinning planes. The deformation process in the low temperature crystallized Ti50 Ni25 Cu25 ribbon involves mainly the detwinning of the (1 1 1) type I twin which is isotropic due to the fiber texture. Such isotropic detwinning results in smaller plateau strain and earlier dislocation process. This further leads to lower shape memory strain and high spring back. The TWME developed in Ti50 Ni25 Cu25 ribbon is related to the specially arranged dislocations produced during martensite deformation. The dislocations would create the directional internal stress field and causes preferential growth of martensite variants [12]. This low TWME should be related to the fiber texture that leads to formation of randomly oriented of martensite variants. In addition, the sample with a smaller grain size usually has a higher yield stress and suppresses the generation of dislocations. This further results in less preferentially oriented martensite variants upon cooling, thus, smaller two-way memory strain. 4. Conclusions 1. The initially amorphous ribbon is fully crystallized under RTA at 400 ◦ C for 30 s. However, under CTA the ribbon remains amorphous even after annealed at 400 ◦ C for 30 min. The crystallization at low temperature under rapid heating is attributed to the release of the extra energy stored in the free-volume of amorphous state.
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2. No precipitates of B11–TiCu phase are found in the Ti50 Ni25 Cu25 ribbon annealed at 400 ◦ C for up to 300 s under RTA. 3. With increasing the annealing time, the transformation temperatures increase. The dependence of transformation behavior on the annealing time could be related to the constraint from the amorphous–crystalline interface and the grain boundaries. 4. For the sample annealed at 400 ◦ C for 300 s, a one-way recovery strain of 2.7% is obtained with the tensile strain of 5.5%. The stress–strain curves show non-flat stress-plateau and high spring back. A TWME is induced by martensite deformation. Both the deformation behavior and TWME are likely related to the isotropic detwinning as a result of formation of randomly oriented martensite variants. The small grain size also contributes to the large spring back and low TWME. References [1] W.J. Moberly, K.N. Melton, in: T.W. Duerig, K.N. Melton, D. St¨okel, C.M. Wayman (Eds.), Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann Ltd., London, UK, 1990, p. 46. [2] Z.L. Xie, J. Van Humbeeck, Y. Liu, L. Dalaey, Scripta Mater. 37 (1997) 363. [3] Yeon-Wook Kim, Tae-Hyun Nam, Scripta Mater. 51 (2004) 653. [4] Y. Liu, Z.L. Xie, Y.X. Tong, C.W. Lim, J. Alloys Compd. 416 (2006) 188. [5] Y. Liu, Mater. Sci. Eng. A 354 (2003) 286.
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