Stable and large superelasticity and elastocaloric effect in nanocrystalline Ti-44Ni-5Cu-1Al (at%) alloy

Accepted Manuscript Stable and large superelasticity and elastocaloric effect in nanocrystalline Ti-44Ni-5Cu-1Al (at%) alloy Hong Chen, Fei Xiao, Xiao Liang, Zhenxing Li, Xuejun Jin, Takashi Fukuda PII:

S1359-6454(18)30627-X

DOI:

10.1016/j.actamat.2018.08.003

Reference:

AM 14755

To appear in:

Acta Materialia

Received Date: 26 June 2018 Revised Date:

30 July 2018

Accepted Date: 1 August 2018

Please cite this article as: H. Chen, F. Xiao, X. Liang, Z. Li, X. Jin, T. Fukuda, Stable and large superelasticity and elastocaloric effect in nanocrystalline Ti-44Ni-5Cu-1Al (at%) alloy, Acta Materialia (2018), doi: 10.1016/j.actamat.2018.08.003. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Graphical Abstract

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Stable and large superelasticity and elastocaloric effect in nanocrystalline Ti-44Ni-5Cu-1Al (at%) alloy

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Thermal-induced B2-B19-B19’ MT

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Annealing at 673K/5min

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Cold rolling at RT

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ACCEPTED MANUSCRIPT Stable and large superelasticity and elastocaloric effect in nanocrystalline Ti-44Ni-5Cu-1Al (at%) alloy Hong Chena, Fei Xiaoa,c∗, Xiao Lianga, Zhenxing Lia, Xuejun Jinb** and Takashi Fukudac∗∗∗ State Key Lab of Metal Matrix Composite, School of Materials Science and

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a

Engineering, Shanghai Jiao Tong University, 800 Dong Chuan Road, Shanghai 200240, P. R. China b

Institute of Advanced Steels and Materials, School of Materials Science and

China

Department of Materials Science and Engineering, Graduate School of Engineering,

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c

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Engineering, Shanghai Jiao Tong University, Shanghai 200240, People’s Republic of

Osaka University, 2-1, Yamada-oka, Suita, Osaka 565-0871, Japan

Abstract

Superelastic behavior and elastocaloric effect were investigated in a Ti-44Ni-5Cu-1Al

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(at%) alloy subjected to various thermomechanical treatments. The specimen heat-treated at 673 K for 5 minutes after hot rolling and subsequent cold rolling exhibited excellent superelastic strain of 4.9% with a small stress hysteresis of

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90 MPa when the maximum tensile stress was 500 MPa. This specimen also exhibited a large elastocaloric effect with a temperature decrease of 17 K when the stress of 500 MPa was removed adiabatically. No remarkable deterioration was observed for

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the superelastic strain and elastocaloric effect up to 5000 mechanical cycles. The maximum superelastic strain obtained was 6.8% under a tensile stress of 750 MPa. Transmission electron microscope observation and in-situ X-ray diffraction analysis under tensile stress revealed that the average grain size of the specimen is about 40 nm, and the specimen exhibits a successive B2-B19-B19’ transformation. ∗

Corresponding author. Tel.: +86 21 54745560; fax: +86 21 34203098. E-mail address: [email protected] (F. Xiao). ** Corresponding author. Tel.: +86 21 54745560; fax: +86 21 34203098. E-mail address: [email protected] (X. Jin). ∗∗∗ Corresponding author. Tel.: +81 6 6879 7483; fax: +81 6 6879 7522. E-mail address: [email protected] (T. Fukuda).

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Keywords: Ti-Ni; shape memory alloys; martensitic transformation; elastocaloric

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effect; fatigue

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ACCEPTED MANUSCRIPT 1. Introduction Superelasticity in shape memory alloys (SMAs) is now widely used for practical applications including biological devices [1-3]. It is caused by stress-induced martensitic transformation (MT) and the reverse transformation. Since MT is

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associated with latent heat, stress-induced MT causes temperature rise or drop of the specimen when the stress is applied or removed adiabatically. This behavior is frequently referred to as elastocaloric effect [4-8]. Recently, elastocaloric effect in SMAs has attracted much attention because it has high potential to be used for new

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refrigeration system [9-11]. For both superelasticity and elastocaloric effect, high fatigue resistance and good workability are important factors for practical

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applications.

Among many SMAs, Ti-Ni SMAs are most widely used because of their excellent mechanical properties and workability [12]. Interestingly, elastocaloric effect is also excellent in Ti-Ni SMAs because of the high latent heat associated with the stress-induced MT [13-15]. It is reported that adiabatic temperature decrease of about

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21 K can be reached by stress removal [16].

Although many studies have been made to improve superelasticity and fatigue properties of the Ti-Ni SMAs [17-19], their improvements are still the most

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challenging subject. Reduction of the grain size is effective to improve mechanical properties. For example, Y. H. Kim et al. [20] and Q. P. Sun et al. [21-23] demonstrated improvement of superelasticity in nanocrystalline Ti-Ni SMAs obtained

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by cold rolling. Still, its fatigue property was not satisfactory [24]. Presumably, inferior lattice compatibility between the parent and martensite phases will be the main reason for the insufficient fatigue property [25, 26]. There are three martensite phases in Ti-Ni SMAs. They are the R-phase, the

B19-phase and the B19’-phase [27]. Although the superelastic strain is the largest for the B2-B19’ transformation, the lattice compatibility between the B2-phase and the B19’-phase is not good for most alloys. Therefore, defects accumulate in the specimen through repeating the B2-B19’ transformation, resulting in insufficient 3

ACCEPTED MANUSCRIPT fatigue property [28-30]. The lattice compatibility is excellent between the B2-phase and the R-phase, and excellent fatigue properties are reported for the B2-R transformation [31]. Nevertheless, the transformation strain and latent heat are small for the B2-R transformation, which limit the use of this transformation [32, 33]. It is

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largely recognized that the superelasticity and fatigue property are intermediate for the B2-B19 transformation [34-36]. Recently, an excellent lattice compatibility between the B2-phase and the B19-phase was reported in a Ti-Ni-Cu-Pd alloy through tuning the composition of alloy [37]. In this type alloy, quite a small transformation

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hysteresis and drastic improvement of fatigue property were demonstrated [38, 39].

Considering the improvement of superelasticity by introducing fine microstructure

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and better fatigue properties for the B2-B19 transformation, we may expect an excellent superelasticity by introducing fine microstructure in alloys exhibiting the B2-B19 transformation. In Ti-Ni SMAs, substitution of Cu for Ni is the most representative method to form the B19 martensite. It is generally recognized that Cu content of 7.5 at% and more is necessary to induce the B19-phase, and higher Cu

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content is preferable for better lattice compatibility between the parent and the martensite phase [26, 40]. However, addition of a large amount of Cu significantly deteriorates workability. In order to obtain fine grain using cold rolling, smaller Cu

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content is preferable.

We considered the Ti-44Ni-5Cu-1Al (at%) could be a candidate to obtain a nano-scale microstructure through cold rolling. We can expect high cold workability

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in this alloy because of small Cu content [41]. In addition, this alloy was reported to exhibit a successive B2-B19-B19’ transformation although the Cu content is less than 7.5 at%. In this alloy, the B19-phase firstly forms in the B2 matrix, then the B19’-phase forms in the B19-phase [42, 43]. It seems that doping of Al effectively stabilize the B19-phase compared to the B19’-phase. Moreover, the transformation hysteresis of this alloy after homogenization heat-treatment is nearly a half of that of a Ti-45Ni-5Cu alloy subjected to the same heat-treatment [42], suggesting that introduction of defects through the MT is small in this alloy. 4

ACCEPTED MANUSCRIPT We found that severe cold working is possible in the Ti-44Ni-5Cu-1Al alloy and nano-scale microstructure is produced through cold rolling and subsequent heat-treatment at 673 K. Moreover, we found the heat-treated alloy shows an

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excellent superelasticity and elastocaloric behavior with negligibly little fatigue.

2. Experiments

An ingot of Ti-44Ni-5Cu-1Al (at%) alloy was prepared by vacuum induction melting in a graphite crucible. The ingot was cut into a slab, and the slab was hot

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rolled at 1123 K to 50% thickness reduction (final thickness ~2 mm). We call this specimen as hot rolled (HR) specimen. Then the hot rolled sheet was further cold

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rolled to 50% thickness reduction by 5 passes (final thickness~1 mm). We term this specimen as cold rolled (CR) specimen. To fabricate different states of the Ti-Ni-Cu-Al alloys, the CR sheets were annealed at 673 K for 5 min and 873 K for 5 min in evacuated quartz tubes then quenched into ice water. These annealed sheets were referred to as CR673 and CR873 specimens, respectively.

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The transformation characteristics of these samples were measured by a NETZSCH model 200F3 differential scanning calorimeter (DSC) with a cooling/heating rate of 10 K/min, and by standard four-point contact electrical resistivity (ER) measurements

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with a cooling/heating rate of 2 K/min.

Dog-bone shaped specimens with gauge length of 20 mm, width of 3 mm, and thickness of 1 mm for tensile tests were cut along the rolled direction by using slow

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wire-electrode cutting method. The specimens were mechanically polished to a mirror-like surface. Tensile tests were performed at different constant temperatures by an Instron-5966 mechanical testing machine with a loading/unloading strain rate of 5×10-3 s-1, and the strain was determined by video extensometer. The stress controlled cycling tests were conducted by a MTS Landmark servohydraulic machine at a frequency of 0.2 Hz. The elastocaloric effect was detected at a fast unloading strain rate of 2 s-1 to achieve nearly adiabatic condition. The temperature change of the specimen was monitored by a T-type thermocouple welded on the surface of the 5

ACCEPTED MANUSCRIPT specimen. The in-situ structural evolution under the tensile stress was evaluated by a PANalytical Empyrean X-ray diffractometer (XRD) with a Cu-Kα radiation source. Samples for transmission electron microscope (TEM) observation were mechanically

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ground to a thickness of about 80 µm, then double jet electro-polished using a 10 vol.% acetic acid and 90 vol.% methanol solution at 283 K. Microstructural observations

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3. Results

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3.1 Martensitic transformation characteristics

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were performed at RT (~293 K) by using a JEOL 2100F at an accelerating voltage of

Figure 1 shows the DSC heat flow in the temperature range of 123-393 K for the specimens with different thermomechanical treatments (HR, CR, CR673 and CR873). Although only one peak can be seen, the successive B2-B19-B19’ MT occurs in all the specimens according to a previous report [42] and the present X-ray experiments

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described below. The reason of single peak is attributed to the overlap of the B2-B19 MT and the B19-B19’ MT. The peak of the CR specimen is very broad and hardly detectable. Such behavior was also observed in binary Ti-Ni alloys after severe cold

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rolling [44]. After annealing at 673 K for 5 min (CR673), the transformation peak became detectable again. This could be due to the grain growth through annealing treatment. The DSC curve of the CR873 specimen is almost the same as that of the

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HR specimen, suggesting that the influence of cold working completely disappears when the specimen is heated to 873 K. The latent heats (∆H) of the three specimens are indicated in Figure 1. ∆H of the

HR and CR873 specimens (~1.2 kJ/mol) indicated the typical first-order MT. ∆H of the CR673 specimen is about 0.5 kJ/mol. This value is less than 1/2 of the latent heat of the HR and CR873 specimens. One possibility is that the latent heat is not correctly evaluated because of the broadness of the peak in the CR673 specimen. Another possibility is that the volume fractions of the B2, B19 and B19’ phases are different 6

ACCEPTED MANUSCRIPT between the HR and CR673 specimens even below the MT finish temperature (Mf). Due to significant broadness of the transformation peak of the CR specimen, it is difficult to determine the latent heat as well as transformation temperatures. Since it is difficult to detect MT temperatures from the DSC curve for the CR

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specimen and the CR673 specimen, we made electrical resistivity measurements. Figure 2 presents temperature dependence of electrical resistivity (ρ-T curve) of each specimen. The ρ-T cooling curve of the HR specimen shows a sharp increase in resistivity at ~306 K due to MT. No clear two step changes in resistivity can be seen

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for both the cooling and heating processes although the specimen was reported to show a successive B2-B19-B19’ transformation [42]. This result is consistent with the

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DSC measurement. Again overlap of the B2-B19 and B19-B19’ transitions will be the reason for such behavior. The ρ-T curves of the CR and CR673 specimens change rather gradually in a wide temperature range. This behavior is consistent with the broad DSC peak observed in Figure 1. The ρ-T curve of the CR873 specimen resembles that of the HR specimen, being again consistent with the DSC result

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(Figure 1).

By employing a tangent method on resistivity curves, we determined the MT temperatures (Ms, Mf, As and Af) and the values are presented in Figure 3(a). The

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thermal hysteresis is also evaluated as (Af - Ms), and the value is shown in Figure 3(b). It is found that Ms and Af change slightly, but Mf and As change significantly by cold rolling and subsequent heat-treatment. A remarkable result is that the CR673

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specimen shows very small temperature hysteresis of ~9 K. The successive thermal-induced B2-B19-B19’ MT in the HR specimen was already

confirmed in a previous work [42]. In order to confirm such successive transformation behavior in the CR673 specimen, we made in-situ XRD in a temperature range of 123-423 K, and the results of the cooling and the subsequent heating processes are shown in Figure 4. The indexes of peaks are given in (b). At 423 K (>Af), the main reflection is 110B2. As temperature decreases to 258 K, 020B19 and 010B19’ reflections appear. On further cooling to 233 K, the intensity of 020B19 and 010B19’ increases 7

ACCEPTED MANUSCRIPT while that of 110B2 decreases. At 123 K, it is almost impossible to distinguish these peaks separately because of the broadening and overlap of these peaks. In the heating process, the peak profile changes reversely. It is obvious from these results that the CR673 specimen exhibits the successive B2-B19-B19’ MT in the cooling process and

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the reverse transformation in the heating process. Incidentally, we notice a shoulder near the position of 101B19’ reflection even at 423 K (>Af). Presumably, a small part of the specimen remains to be the B19’-phase up to 423 K because of residual stress in

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the CR673 specimen.

3.2 Microstructures

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The microstructure of each specimen was analyzed by TEM at RT. Figure 5(a) shows the bright field images of the HR specimen. The grain size of the HR specimen is in micro-scale. The selected area diffraction pattern (SADP) taken from regions of B, C and D are presented in (b)-(d), respectively. The diffraction patterns (b), (c) and (d) are indexed with the B2-, B19- and B19’-phases, respectively. We notice fine twin

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in the B19’-phase. This is characteristic to the compound twin formed by the B19-B19’ transformation. This result again supports that the HR specimen exhibits successive B2-B19-B19’ MT as reported before [42].

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The bright field image and the corresponding diffraction pattern of the CR specimen are shown in Figure 6. Typical elongated microstructure along the rolled direction (indicated by the red arrow) is observed in the bright field image. Broad

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diffraction ring of the diffraction pattern suggests the existence of an amorphous structure in the CR specimen. The average grain size was determined to be about 8 nm from dark field images (not shown here). The typical microstructure of the CR specimen is consistent with the previous reports of cold rolled Ti-Ni binary SMAs [44]. Figure 7(a) shows the bright field image and the corresponding diffraction pattern of the CR673 specimen. The average grain size is ~40 nm, and the diffraction rings and spots are sharper than those of the CR specimen. The diffraction pattern is 8

ACCEPTED MANUSCRIPT essentially indexed by the parent phase with the B2-type structure, but there are additional weak reflections as indicated by red arrows in Figure 7(a). In order to identify the origin of these weak reflections, we made high resolution TEM observation, and the results are shown in (b)-(d). Fast Fourier transformation of the

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red rectangular areas show the pattern of the cubic B2 parent phase (b’), the orthorhombic B19 martensite phase (c’) and the monoclinic B19’ martensite phase (d’). This implies the CR673 specimen undergoes a successive B2-B19-B19’ MT like the HR specimen.

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Figure 8 shows microstructure of the CR873 specimen. The average grain size is ~1 µm, which is much larger than the CR673 specimen but apparently smaller than

3.3 Mechanical behaviors

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the HR specimen.

Typical stress-strain curves for the HR (a), CR (b), CR673 (c) and CR873 (d) specimens measured above Af temperature are shown in Figure 9. The stress-strain

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curve of the HR specimen (a) at 332 K (Af +10 K) shows a typical plateau caused by the stress-induced MT. The critical stress for inducing MT (σc) is ~410 MPa, which is determined by the tangent method as indicated in the figure. Although a part of the

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strain recovers, a large part of the strain (~67%) is unrecoverable in the stress removing process. Similar results were obtained at higher temperatures. In addition, most of the residual strain did not recover by heating the specimen after removing the

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stress. These results imply that the unrecoverable strain is not due to a large stress hysteresis but due to plastic deformation of the specimen. This poor superelasticity of the HR specimen is consistent with a previous report [45]. The stress-strain curve of the CR specimen at 335 K (Af +10 K) does not show a

clear plateau like the HR specimen as shown in Figure 9(b). Nevertheless, it is not linear, it is slightly convex upwards. In addition, we notice a small hysteresis between the loading and unloading processes. This implies that a small part of the specimen exhibits stress-induced MT. The maximum reversible strain is ~2.5%, which is 9

ACCEPTED MANUSCRIPT induced under a tensile stress of ~880 MPa. Unrecoverable strain was introduced when the stress of 1000 MPa was applied as indicated in Figure 9 (b). In the stress-strain curve of the CR673 specimen at 318 K (Af +8 K), we notice an oblique plateau characteristic to stress-induced MT but the strain caused by the

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plateau (~4.5%) is smaller than that of the HR specimen. We consider this strain as the transformation strain εT. The recoverable strain reaches ~6.8% under a maximum stress of ~800 MPa. The stress hysteresis is ~150 MPa, which is much smaller than that of the HR specimen. The value of σc at 318 K is determined to be ~260 MPa as

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indicated in the figure. Similar results were obtained at different test temperatures (Figure s1); and the value of σc is plotted as a function of temperature in Figure 10. It

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increases linearly as test temperature increases.

The superelasticity of the CR873 specimen at 325 K (Af +10 K) is apparently better than that of the HR specimen, but is not as good as that of the CR673 specimen as shown in (d). A large residual strain of ~1% appears even when the maximum strain

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of below 5% is applied at 325 K (Af +10 K).

3.4 In-situ structural evolution under the tensile stress In order to confirm the stress-induced MT during the tensile test, in-situ XRD

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measurement was conducted at RT, and the structural evolution during the loading and unloading processes for the HR (a, a’), CR (b, b’), CR673 (c, c’) and CR873 (d, d’) specimens are present in Figure 11. The peaks are indexed in (a).

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The X-ray profile of the HR specimen (a) is composed of the B2-phase, B19-phase

and B19’-phase under zero stress, which is consistent with the TEM result (Figure 5). As the tensile stress increases, the intensities of 110B2 and 101B19’ peaks decrease, while those of the 020B19 and 111B19’ peaks increase. These peaks gradually merged together with increasing stress. The diffraction pattern was almost unchanged during the unloading process. This means the structure is unrecoverable during the unloading process for the HR specimen. This implies that the unrecoverable strain in Figure 9(a) is mainly caused by irreversible nature of the stress-induced MT in the HR specimen. 10

ACCEPTED MANUSCRIPT The X-ray profile of the CR specimen (b) shows a broad peak with its center near 110B2. This broad peak includes peak positions for 020B19 and 111B19’. The peak position shifts slightly toward high angle with increasing the tensile stress, and returns back to the original state after removing the applied stress. This behavior corresponds

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to the non-linear deformation of the CR specimen (Figure 9(b)). The structural evolution for the CR673 specimen during the loading and unloading processes is reversible as shown in Figures 11(c) and (c’). The initial diffraction pattern for the CR673 specimen resembles that of the HR specimen, but with

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relatively broader peaks. Due to the broadness of the peaks, it is difficult to separate the 020B19 and 111B19’ in the CR673 specimen. But the reversible decrease and

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increase for the pre-existed 101B19’ reflection can be clearly observed during the loading and unloading processes as indicated by the red arrows. This reversible structural change for the pre-existed martensite phase will be further described latter. The successive stress-induced B2-B19-B19’ transformation can be clearly observed in the CR873 specimen, but the reversible structural evolution cannot be seen shown

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in Figures 9(d) and (d’). Notice that the peaks are broader than those of the HR specimen, which could be due to the remained internal stress caused by cold rolling

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process in the CR873 specimen.

3.5 Fatigue behavior and elastocaloric effect Both the temperature hysteresis and stress hysteresis are small for the successive

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B2-B19-B19’ transformation in the CR673 specimen compared with the HR specimen and the CR873 specimen as shown in Figures 2 and 9. Therefore, we may expect a good fatigue behavior in the CR673 specimen. Thus, a long-term loading/unloading test under the maximum tensile stress of ~500 MPa was performed on the CR673 specimen at 318 K (>Af+8 K). The strain induced by this stress is 4.9%. We also made a comparison test on the B2-B19’ transformation in a Ti-50.8Ni (at%) binary alloy, which was prepared by the same cold rolling and subsequent heat-treatment (termed as Ti-Ni CR673). To obtain the same reversible strain of 4.9%, the maximum tensile 11

ACCEPTED MANUSCRIPT stress of ~580 MPa was applied at 338 K (>Af +10 K) for the Ti-Ni CR673 specimen. The stress-strain curves for the CR673 specimen is shown in Figure 12(a) and that of Ti-Ni CR673 specimen is shown in Figure 12(b). The stress-strain curves of the CR673 specimen hardly changed through 5000

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mechanical cycles. However, that for the Ti-Ni CR673 specimen obviously changed even by 10 mechanical cycles, where the maximum reversible strain and the σc both decreased with increasing cycles as reported in previous papers [18, 46]. Notice that the stress hysteresis for the CR673 specimen (~90 MPa) is much smaller than that of

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the Ti-Ni CR673 specimen (~270 MPa). The Ti-Ni CR673 specimen broke within several hundred stress cycles, which was also reported in other reference [24].

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Due to the stable mechanical behavior for the CR673 specimen of the present Ti-44Ni-5Cu-1Al alloy, it is expected that the elastocaloric effect could also be stable. In order to know the stability of the elastocaloric effect, we measured the temperature changes for both the CR673 and Ti-Ni CR673 specimens during the quick unloading process and the results are presented in Figures 13(a) and (b), respectively. The stress

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was increased at a relatively low strain rate of 10-3 s-1 up to ~500 MPa for the CR673 specimen and ~580 MPa for the Ti-Ni CR673 specimen then it was immediately removed at a high strain rate of 2 s-1 to realize adiabatic condition. The temperature

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gradually increased during the loading process and sharply decreased during the unloading process. The temperature decrease in the stress removing process is ~17.4 K for the CR673 specimen and ~25.6 K for the Ti-Ni CR673 specimen. As is

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expected, the decay of elastocaloric effect is small for the CR673 specimen up to 5000 cycles. On the other hand, a significant decay in the caloric effect can be seen through only 10 mechanical cycles in the Ti-Ni CR673 specimen.

4. Discussion In most Ti-Ni binary SMAs, the large reversible strain caused by the B2-B19’ MT significantly decay by several mechanical cycles [28, 30, 31]. In the present work, through introducing the B19 martensite phase and the nanocrystalline grains, the 12

ACCEPTED MANUSCRIPT CR673 specimen showed stable mechanical behavior and elastocaloric effect (Figures 12 and 13). We discuss the mechanism below.

4.1 Effects of cold rolling and annealing on superelasticity

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The unrecoverable strain of the HR specimen is attributed to the relatively low strength of the HR matrix, where the irreversible plastic deformation is introduced at relatively low stress level. The refinement of grain size through cold rolling effectively strengthens the matrix of the Ti-44Ni-5Cu-1Al alloy. The yield stress

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exceeds 880 MPa in the CR specimen as shown in Figure 9(b). But the CR specimen includes high density of defects and an amorphous structure which significantly

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suppress the MT [44]. Therefore, the DSC peak caused by thermally-induced MT (Figure 1) and the stress-plateau caused by stress-induced MT (Figure 9(b)) are both unclear in the CR specimen. Thus the superelasticity is small in the CR specimen. By employing an annealing treatment, the MT in the CR specimen is resumed because of recrystallization and annealing out of a part of defects. However, the

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annealing treatment also promotes grain growth which eliminates the strengthening effect of fine grain as seen in the CR873 specimen. When the annealing temperature is 673 K, the density of the defects decreases to a certain degree while retaining the fine

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microstructure with the average grain size of ~40 nm (Figure 7). Then MT occurs while keeping the strengthening effect of the matrix.

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4.2 Diffuse B2-B19-B19’ transformation in the CR673 specimen As seen in the DSC curve and electrical resistivity curve, the thermally-induced

successive B2-B19-B19’ MT occur in a wide temperature range in the CR673 specimen. Also, as seen in the stress-strain curve, the stress-induced B2-B19-B19’ MT occur in a wide stress range. Such a diffuse nature of the MT could be another reason for the stable superelasticity and elastocaloric effect in the CR673 specimen. There are several characteristic features for the diffuse MT in the CR673 specimen. First, the X-ray profile of the martensite phase is very broad as seen in Figure 4. 13

ACCEPTED MANUSCRIPT Presumably, the lattice parameters of the martensite phase distribute continuously in the specimen. Second, the peak of the DSC curve is relatively broad and the latent heat evaluated from the DSC curve is obviously smaller than that of the HR specimen. Third, the transformation strain is smaller than that observed for the sharp MT in the

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HR specimen. Finally and most importantly, the stress hysteresis of the MT is much smaller than that in the HR specimen. These features will be the consequence of lattice compatibility between the parent and martensite phases.

The small grain size (~40 nm) of the CR673 specimen will be responsible for these

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characteristic features. The growth of the martensite phase would be restricted to nano-scaled grains [30, 47]. This enforces the whole transformation as a diffuse one.

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Defects introduced during the cold rolling process in the matrix will be also responsible for the diffuse nature of the transformation. These defects will act as nucleation cites which reduce the potential barrier of transformation, resulting in small hysteresis. They also create inhomogeneous internal stress field which prevents sharp MT, resulting in diffuse transformation.

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Meanwhile, the relative low stress hysteresis (~90 MPa) and σc (~260 MPa) in the Ti-44Ni-5Cu-1Al CR673 specimen compared with that for the Ti-Ni CR673 specimen (~270 MPa and ~510 MPa) implies the improvement of the lattice compatibility

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through introducing the successive B2-B19-B19’ transformation (in-situ XRD result in Figures 11(c) and (c’)). This means the successive B2-B19-B19’ transformation also contributes to the stable mechanical behavior in the Ti-44Ni-5Cu-1Al CR673

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specimen.

In addition, the stable superelasticity could also be related to the residual martensite

phase which exists up to 423 K (Figure 4). It is speculated that a part of superelasticity could be caused by reversible reorientation of the B19’ martensite phase like rubber-like deformation [48]. The increase and decrease of the 110B19’ reflection seen in Figure 11 could be caused by the reorientation of the B19’ martensite variants. The residual martensite in the CR673 specimen could be stabilized by the internal stress introduced by cold rolling. 14

ACCEPTED MANUSCRIPT Incidentally, we are curious about whether the Clausius-Clapeyron equation ௗఙ

(

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= −Δܵ/Δߝ் ) can be applied for the CR673 specimen in which MT occurs in wide

temperature and stress ranges. From Figure 10, the slope (dσ/dT) is calculated to be ~5.4 MPa/K. The transformation strain ΔεT is estimated to be 4.5% at 318 K as

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mentioned before. Then the entropy change Δܵ is evaluated to be ~ -2.0 J/(mol·K). On the other hand, from the latent heat of the CR673 specimen, the entropy change is evaluated to be Δܵ = ∆‫ܪ‬/ܶ଴ = -1.8 J/(mol·K). The agreement of the entropy change evaluated by two methods implies that we may use the Clausius-Clapeyron equation

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even for the diffuse MT of the CR673 specimen. Nevertheless, we have to be aware that the entropy change evaluated here is not the one for the complete B2-B19-B19’

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transformation, but for a partial transformation.

The energy efficiency of materials (η) is an important parameter for cooling technologies. According to Ref. [49], η is defined as η = |Q/W|, where mechanical work, W, is done to drive highly reversible caloric effects in an isothermal body, whose entropy is thus modified such that heat, Q, flows to or from the thermal bath.

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The value of |Q| was evaluated as |–c∆T| [9], where the c is specific heat capacity and ∆T is adiabatic temperature change. The value of |W| was evaluated as the enclosed area of the stress-strain curve [9]. The values of |Q|, |W| and η thus evaluated were

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51.2 J/cm3, 4.6 J/cm3 and ~11 for the Ti-44Ni-5Cu-1Al CR673 specimen while they were 58.8 J/cm3, 9.5 J/cm3 and ~6 for the Ti-Ni CR673 specimen. The value η=11 in

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the former specimen is higher than that of any mechanocaloric material reported in [49] (2.7~8.4).

5. Conclusions

The transformation characteristics, microstructures, mechanical behaviors, and in-situ structural evolution were systematically investigated in a Ti-44Ni-5Cu-1Al (at%) alloy subjected to various thermomechanical treatments: hot rolled (HR), cold rolled (CR), cold rolled followed by aging at 673 K for 5min (CR673), and cold rolled followed by aging at 873 K for 5min (CR873) specimens. The key conclusions are as 15

ACCEPTED MANUSCRIPT follows: (1) The HR and CR873 specimens have micro-scaled grains and show the stress-induced typical first-order B2-B19-B19’ MT during the loading process, but they are irreversible during the unloading process.

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(2) The thermal- and stress-induced MTs is partially suppressed in the CR specimen due to the introduction of dense defects and amorphous phase introduced by cold rolling. The CR specimen shows a non-linear superelasticity.

(3) The CR673 specimen with nano-scaled grains (~40 nm) shows a diffuse

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B2-B19-B19’ MT. A large reversible strain ~6.8% with relative low σc (~260 MPa) and stress hysteresis (~150 MPa) is obtained under a tensile stress of 750 MPa.

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(4) The CR673 specimen shows a stable superelastic strain of 4.9% with a small stress hysteresis of 90 MPa and stable elastocaloric effect (∆T~17.4 K) under a tensile stress of ~500 MPa. Deterioration of superelasticity and elastocaloric effect is negligibly small up to 5000 mechanical cycles. High value of energy efficiency

Acknowledgement

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(η ~11) was obtained in the Ti-44Ni-5Cu-1Al CR673 specimens.

This work was funded by the National Youth Science Foundation (No. 51501113), the Key

R&D

Program

of

China

(No.

2017YFB0703003,

No.

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National

2017YFB0406000) and National Natura Science Foundation of China (U1564203). The authors would like to thank Prof. Q.P. Sun (The Hong Kong University of

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Science and Technology) for his valuable support.

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ACCEPTED MANUSCRIPT Figure captions Figure 1. DSC cooling and heating curves of the Ti-44Ni-5Cu-1Al (at%) alloy measured at a rate of 10 K/min after different thermomechanical treatments (HR, CR,

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CR673, CR873). The latent heats are indicated in the figure.

Figure 2. Temperature dependence of the electrical resistivity of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments (HR, CR, CR673, CR873).

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Figure 3. (a) The martensite transformation temperatures (Ms, Mf, As and Af) of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments (HR, CR,

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CR673, CR873). (b) The thermal hysteresis evaluated as (Af - Ms).

Figure 4. In-situ X-ray diffraction profiles of Ti-44Ni-5Cu-1Al (at%) CR673 specimen at different testing temperatures during the cooling (a) and heating (b)

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processes.

Figure 5. (a) Bright field image for the Ti-44Ni-5Cu-1Al (at%) HR specimen. (b), (c), (d) Selected area diffraction patterns taken from the regions marked B, C and D,

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respectively.

Figure 6. Bright field image and corresponding diffraction pattern for the

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Ti-44Ni-5Cu-1Al (at%) CR specimen. Arrow indicates the rolling direction.

Figure 7. (a) Bright field image and selected area diffraction pattern for the Ti-44Ni-5Cu-1Al (at%) CR673 specimen. (b), (c), (d) High resolution TEM images and (b’), (c’), (d’) corresponding fast Fourier transformation for the red rectangular areas, respectively.

Figure 8. Bright field image for the Ti-44Ni-5Cu-1Al (at%) CR873 specimen. 22

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Figure 9. The stress-strain curves of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments. (a) HR specimen tested at 332 K, (b) CR specimen tested at 335 K, (c) CR673 specimen tested at 318 K, (d) CR873 specimen tested at

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325 K. The critical stresses for inducing martensite transformation are indicated in the figure.

Figure 10. The temperature dependence of the critical stress for inducing martensite

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transformation for the Ti-44Ni-5Cu-1Al (at%) CR673 specimen.

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Figure 11. In-situ X-ray diffraction profiles of Ti-44Ni-5Cu-1Al (at%) alloys at different tensile stresses during loading and unloading processes for the HR (a, a’), CR (b, b’), CR673 (c, c’) and CR873 (d, d’) specimens, respectively.

Figure 12. The stress-strain curves responses during stress-controlled fatigue tests of

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different specimens: (a) the Ti-44Ni-5Cu-1Al (at%) CR673 specimen under the maximum tensile stress ~500 MPa at 318 K, (b) the Ti-50.8Ni (at%) CR673 specimen under the maximum tensile stress ~580 MPa at 338 K. The maximum reversible

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strains at the initial state for these two specimens are close to each other of ~4.9%.

Figure 13. The adiabatic temperature change, ∆T, as a function of time in different

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stage, (a) 1st, 1500th and 5000th for the Ti-44Ni-5Cu-1Al (at%) CR673 specimen under the maximum tensile stress ~500 MPa, (b) 1st, 2nd and 10th for the Ti-50.8Ni (at%) CR673 specimen under the maximum tensile stress ~580 MPa. The ∆T caused by the reverse martensite transformation are indicated in the figure.

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Figure 1. DSC cooling and heating curves of the Ti-44Ni-5Cu-1Al (at%) alloy measured at a rate of 10 K/min after different thermomechanical treatments (HR, CR, CR673, CR873). The latent heats are indicated in the figure.

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Temperture (K) Figure 2. Temperature dependence of the electrical resistivity of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments (HR, CR, CR673, CR873).

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Figure 3. (a) The martensite transformation temperatures (Ms, Mf, As and Af) of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments (HR, CR, CR673, CR873). (b) The thermal hysteresis evaluated as (Af - Ms).

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CR673 - Cooling

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Figure 4. In-situ X-ray diffraction profiles of Ti-44Ni-5Cu-1Al (at%) CR673 specimen at different testing temperatures during the cooling (a) and heating (b) processes.

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Figure 5. (a) Bright field image for the Ti-44Ni-5Cu-1Al (at%) HR specimen. (b), (c), (d) Selected area diffraction patterns taken from the regions marked B, C and D, respectively.

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Figure 6. Bright field image and corresponding diffraction pattern for the Ti44Ni-5Cu-1Al (at%) CR specimen. Arrow indicates the rolling direction.

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Figure 7. (a) Bright field image and selected area diffraction pattern for the Ti-44Ni-5Cu1Al (at%) CR673 specimen. (b), (c), (d) High resolution TEM images and (b’), (c’), (d’) corresponding fast Fourier transformation for the red rectangular areas, respectively.

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Figure 8. Bright field image for the Ti-44Ni-5Cu-1Al (at%) CR873 specimen.

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Figure 9. The stress-strain curves of the Ti-44Ni-5Cu-1Al (at%) alloy after different thermomechanical treatments. (a) HR specimen tested at 332 K, (b) CR specimen tested at 335 K, (c) CR673 specimen tested at 318 K, (d) CR873 specimen tested at 325 K. The critical stresses for inducing martensite transformation are indicated in the figure.

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Figure 10. The temperature dependence of the critical stress for inducing martensite transformation for the Ti-44Ni-5Cu-1Al (at%) CR673 specimen.

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Figure 11. In-situ X-ray diffraction profiles of Ti-44Ni-5Cu-1Al (at%) alloys at different tensile stresses during loading and unloading processes for the HR (a, a’), CR (b, b’), CR673 (c, c’) and CR873 (d, d’) specimens, respectively.

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Figure 12. The stress-strain curves responses during stress-controlled fatigue tests of different specimens: (a) the Ti-44Ni5Cu-1Al (at%) CR673 specimen under the maximum tensile stress ~500 MPa at 318 K, (b) the Ti-50.8Ni (at%) CR673 specimen under the maximum tensile stress ~580 MPa at 338 K. The maximum reversible strains at the initial state for these two specimens are close to each other of ~4.9%.

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Figure 13. The adiabatic temperature change, ∆T, as a function of time in different stage, (a) 1st, 1500th and 5000th for the Ti-44Ni-5Cu-1Al (at%) CR673 specimen under the maximum tensile stress ~500 MPa, (b) 1st, 2nd and 10th for the Ti50.8Ni (at%) CR673 specimen under the maximum tensile stress ~580 MPa. The ∆T caused by the reverse martensite transformation are indicated in the figure.