Effect of Heat Treatment Temperature on Martensitic Transformation and Superelasticity of the Ti49Ni51 Shape Memory Alloy

The martensitic transformation and superelasticity of Ti49Ni51 shape memory alloy heat-treatment at different temperatures were investigated. The experimental results show that the microstructures of as-cast and heat-treated (723 K) Ni-rich Ti49Ni51 samples prepared by rapidly-solidified technology are composed of B2 TiNi phase, and Ti3Ni4 and Ti2Ni phases; the microstructures of heat-treated Ti49Ni51 samples at 773 and 823 K are composed of B2 TiNi phase, and of B2 TiNi and Ti2Ni phases, respectively. The martensitic transformation of as-cast Ti49Ni51 alloy is three-stage, A→R→M1 and R→M2 transformation during cooling, and two-stage, M→R→A transformation during heating. The transformations of the heat-treated Ti49Ni51 samples at 723 and 823 K are the A↔R↔M/A↔M transformation during cooling/heating, respectively. For the heat-treated alloy at 773 K, the transformations are the A→R/M→R→A during cooling/heating, respectively. For the heat-treated alloy at 773 K, only a small thermal hysteresis is suitable for sensor devices. The stable σmax values of 723 and 773 K heat-treated samples with a large Wd value exhibit high safety in application. The 773 and 823 K heat-treated samples have large stable strain–energy densities, and are a good superelastic alloy. The experimental data obtained provide a valuable reference for the industrial application of rapidly-solidified casting and heat-treated Ti49Ni51 alloy.


Introduction
The Ni-rich Ti-Ni-based alloys are favored by researchers of engineering materials, functional materials, and biomedical alloys, based on high strength, low density, superelasticity, and shape memory properties [1][2][3][4][5]. These excellent properties enable Ti-Ni-based shape memory alloys (SMAs) to be a functional material which integrates sensing and driving [6,7]. The applications of Ti-Ni-based SMAs mainly include microelectromechanical systems and novel medical devices [6,7]. In microelectromechanical systems, the functional microgrippers, fluid-flow valves, and micromirror actuators are fabricated by using the Ti-Ni-based alloy; in the medical field, the Ti-Ni-based alloy can be successfully used in arterial stents and surgical tools [6,7]. In fact, Ti-Ni-based alloys can be used in different fields because of their shape memory or superelastic properties. SMA with martensitic phase has good shape memory effect, that is, the alloy with a specific composition can be loaded and deformed in martensitic phase state, and retain its deformed shape after unloading, and then the deformed alloy can be restored to the shape before loading when heated to the parent phase state [8]. SMA with the parent phase state has good superelasticity, that is, the strain produced by a specific component alloy under the action of external force exceeds its elastic limit variable, after unloading, the alloy can spontaneously restore to its original shape [9][10][11], and its stress-strain curve presents as non-linear. In addition, the Ti-Ni-based SMA has high damping properties, biocompatibility, and corrosion resistance, and has been applied in aerospace, machinery manufacturing, transportation, civil construction, energy engineering, biomedicine, and daily life [12][13][14][15].
To apply Ti-Ni-based SMAs in different environments, it is necessary to improve the properties of Ti-Ni alloys and to consider preparation methods. Researchers reported that two methods can improve the properties of Ti-Ni SMAs; one is heat treatment [16][17][18][19][20], and the other is the addition of minor elements [13][14][15][18][19][20][21][22]. For the single-crystal Ni-50.9 at% Ti alloy, the Ti 3 Ni 4 phases with the different sizes are precipitated in B2 matrix phase by heat treatment [17]. The reported results show that the room-temperature hardness and recoverable energy of Ti-Ni are a maximum for very small precipitate sizes (10 nm), decrease for intermediate precipitate sizes, and increase for large precipitate sizes [17]. In addition, the Ti 3 Ni 4 precipitate (10 nm) improves the fatigue resistance of the Ti-Ni alloy, compared to the other heat treatments [16]; samples with large incoherent Ti 3 Ni 4 precipitates (500 nm) consistently show significant dislocation activity due to mechanical cycling [16]. The reported Ti 44.5 Ni 44.5 Nb 9 Co 2 alloy annealed at 550 • C shows the highest yield strength, a large complete recovery strain, and high recovery stress, based on the addition of Nb and Co elements [19]. For the aged Ni-rich NiTi alloys, the formation of precipitates increase their strength and alter transformation temperatures [23]. Precipitation formation is an effective method to adjust the shape memory properties of Ti-Ni-based alloys [24]. The chemical composition, volume fraction, and coherency of the formed precipitates in Ni-rich Ti-Ni alloys can influence the transformation temperatures, matrix strength, hardness, and martensite morphology [25][26][27][28]. Thus, combination of aging and precipitates makes it possible to obtain good memory behavior and hyperelastic behavior of Ni-rich Ti-Ni alloys.
In the reported preparation of Ti-Ni alloys, the shape of Ti-Ni alloys is mainly filamentous [20,[29][30][31][32]. The filamentous alloys are mainly melted in a vacuum induction furnace and graphite crucible [31,32]. The ingots are made into wires by pressure forging, rotary forging, drawing, and other processes [31,32]. The obtained wires are then heat-treated to obtain materials used in different environments. To study the mechanical response of crystallographic orientation, some researchers have prepared single-crystal Ti-Ni alloys [16][17][18]28]. Single crystals were grown by Bridgman Technology in an inert gas atmosphere. Different temperatures were used to heat-treated single crystals, and the corresponding effects of different sizes of second phase precipitated on mechanical properties were studied [16][17][18]28]. Powder metallurgy technology can produce large quantities of Ti-Ni ingots with high utilization rate of raw materials; the prepared alloys are then forged and aged to meet the needs of samples in different environments [9,33]. In addition, mechanical alloying is also a commonly used method to prepare the Ti-Ni alloy [9,34]. Ball milling time has an effect on the phase transformation and mechanical properties of the alloy [9,34]. The samples were heat treated again, and different phases were precipitated [9,33,34]. The phase transformation temperature was obtained, so that excellent mechanical properties obtained could meet the needs of engineering applications [9,33,34]. Ti-Ni alloys prepared by the above methods have many processes, or cannot be produced in batches, which increases the processing cost of Ti-Ni alloys. In fact, the preparation of Ti-Ni alloys by water-cooled copper mold casting is rarely reported. The Ti-Ni alloys prepared by rapid solidification can be molded once for some parts with a simple shape, thus shortening the preparation time and cost, and aging treatment of the prepared parts can be carried out to obtain the corresponding transformation temperature and mechanical response based on the application conditions. In this paper, the Ni-rich Ti 49 Ni 51 alloy rods were prepared by rapid solidification technology, and the as-cast rods were heat-treated at different temperatures to gain the different contents and sizes of precipitates, which can affect the transition temperature and superelasticity of the alloy. The effect of heat treatment temperature on martensitic transformation and microstructure was investigated, and the superelastic behavior under cyclic loading was also discussed. Through these studies, the experimental data obtained provide a valuable reference for the industrial application of the rapidly-solidified casting and heat-treated Ti 49 Ni 51 alloy.

Experimental Procedure
Various combinations of pure Ti and Ni (purity of 99.9% or higher) were used to prepare ingot of alloy with nominal compositions of Ti 49 Ni 51 (at%) under a high vacuum (3 × 10 −3 Pa) using a magnetically-controlled tungsten arc-melting furnace (Shenyang Scientific Instruments Co., Ltd. of Chinese Academy of Sciences, Shenyang, China) in an argon atmosphere. These raw materials were obtained from Beijing Xing rong yuan Technology Co., Ltd. (Beijing, China). The raw materials Ti and Ni were in bulk, and the range of sizes was from 5 × 5 × 5 mm 3 to 20 × 20 × 20 mm 3 . The calculated quantities of raw materials were weighed using an analytical balance with a precision of 0.1 mg. After mixing the raw materials, a total of 30 g of raw materials were put into a water-cooled crucible to smelt into alloy ingots. To ensure the uniformity of the chemical composition of the ingots, the ingot was smelted more than 4 times. The melted ingot was placed in a copper crucible for suction casting, melted again, and then sucked into a water-cooled copper abrasives under the action of pressure difference, forming a cylindrical rod with a diameter of 3 mm and a length of 50 mm. Based on expression of cooling rate · T(K/s) = 10/R 2 , (R is diameter, cm) [35], the cooling rate of a rod with a diameter of 3 mm is 111 K·s −1 . The heat treatment temperatures were set to 723, 773, and 823 K, and the heating rate was 20 K min −1 . When heated to the set temperature, the holding time was 30 min. Then the heat-treated sample was cooled to room temperature by furnace.
The samples for mechanical and structural analysis were cut from the as-cast and heat-treated samples by using a slow steel saw. The surfaces of thin sheets with diameter of 3 mm were polished using standard metallographic procedures, consisting of grinding up to 2000 grit with SiC paper and polishing with a colloidal silica suspension. For observing the microstructure, the surfaces of polished samples need to be etched. This was accomplished with a mixed solution of HF, HNO 3 , and H 2 O, and the corresponding volume ratio was 1:4.5:4.5. The microstructures of the prepared samples were observed by optical microscope (OM, Shanghai Changfang Optical Instrument Co., LTD, Shanghai, China). The phases of samples with a diameter of 3 mm were analyzed using an X-ray diffractometer (XRD, Rigaku Company, Tokyo, Japan), at an operating voltage of 30 kV, using Cu-Kα radiation. The transformation temperatures of the samples weighing 20-50 mg were measured by differential scanning calorimetry (DSC) with heating and cooling rates of 0.33 K·s −1 , using a Perkin Elmer DSC 7 instrument (PerkinElmer, Waltham, MA, USA). The cylindrical Ti-Ni samples with a diameter of 3.0 mm and height of~5.0 mm were prepared by subjecting them to uniaxial compression testing at room temperature. The uniaxial compression testing was performed at room temperature using a CMT5105 electronic testing machine (Metis Industrial Systems (China) Co., Ltd., Shanghai, China), at a strain rate of 2.5 × 10 −4 s −1 . The quantitative data were all measured for three times, and the reported data are the average values. The cycle loading and unloading experiments were also carried out on this instrument at a strain rate of 2.5 × 10 −4 s −1 . Figure 1 shows the XRD patterns of as-cast and heat-treated Ti 49 Ni 51 samples. The results show that the microstructures of as-cast and heat-treated Ti 49 Ni 51 samples at 723 K are composed of B2 TiNi phase (CsCl structure), Ti 3 Ni 4 phase (Rhombohedral structure), and Ti 2 Ni phase (Face-centered cubic structure). However, the microstructures of the heat-treated Ti 49 Ni 51 samples at 773 and 823 K are composed of B2-type TiNi phase, and of B2-type TiNi phase and Ti 2 Ni phase, respectively; the diffraction peaks of Ti 3 Ni 4 phase are not observed in Figure 1a. The weak relative diffraction intensity of the Ti 3 Ni 4 phase indicates the low content of Ti 3 Ni 4 phase. Accordingly, the content of the Ti 3 Ni 4 phase of heat-treated samples is lower than that of the as-cast alloy. To clearly see the diffraction peak of Ti 2 Ni phase, Figure 1b is a local enlargement of Figure 1a. The Ti 2 Ni diffraction peak of as-cast alloys is mainly below 30 • . For the 723 K heat-treated alloy, the diffraction peak of Ti 2 Ni phase is 39.6 • ; for the 773 K heat-treated alloy, Ti 2 Ni and Ti 3 Ni 4 phases are not found in XRD patterns. When the heat treatment temperature is further increased, a small amount of Ti 2 Ni precipitates from the 823 K heat-treated alloy. Because the diffraction angle of the strongest diffraction of Ti 2 Ni phase is 41.68 • , the diffraction peaks of the TiNi phase in 723 and 823 K heat-treated alloys become wider, as shown in Figure 1b. Because the intensity of the main diffraction peak of Ti 2 Ni phase is weak, the results show that the content of the Ti 2 Ni phase is relatively small. The cause of Ti 2 Ni phase precipitation needs to be further studied in future work. heat-treated alloys become wider, as shown in Figure 1b. Because the intensity of the main diffraction peak of Ti2Ni phase is weak, the results show that the content of the Ti2Ni phase is relatively small. The cause of Ti2Ni phase precipitation needs to be further studied in future work.  Figure 2 shows the microstructure of as-cast and heat-treated Ti49Ni51 samples. The results show that a large number of Ti3Ni4 phases precipitate along the grain boundary of B2 TiNi phases, and a small amount of Ti3Ni4 phases precipitate inside B2 TiNi phases. Because of the small content of Ti2Ni phase, it is difficult to distinguish Ti2Ni from Ti3Ni4 particles in Figure 2. Therefore, in discussing Figure 2, the precipitation of Ti2Ni phase in TiNi phase is neglected. In addition, the content of Ti3Ni4 phase for the as-cast Ti49Ni51 alloy (see Figure 2a) is significantly higher than that for the heat-treated Ti49Ni51 alloy (see Figure 2b-d). In addition, the content and size of Ti3Ni4 phase decrease with the increase of heat treatment temperature, as shown in Figure 2b-d. The grain sizes of Ti3Ni4 phase for the as-cast Ti49Ni51 alloy (see Figure 2a) are also larger than those of the heat-treated Ti49Ni51 alloy (see Figure 2b-d). In fact, the range of size of Ti3Ni4 particles in the as-cast alloy is 0.2-7 µm; for the 723 K heat-treated alloys, the range of size of Ti3Ni4 particles is 0.2-3 µm; for the 773 and 823 K heat-treated alloys, the ranges of size of precipitates are 0.2-2 µm and 0.2-1 µm, respectively.  heat-treated alloys become wider, as shown in Figure 1b. Because the intensity of the main diffraction peak of Ti2Ni phase is weak, the results show that the content of the Ti2Ni phase is relatively small. The cause of Ti2Ni phase precipitation needs to be further studied in future work.  Figure 2 shows the microstructure of as-cast and heat-treated Ti49Ni51 samples. The results show that a large number of Ti3Ni4 phases precipitate along the grain boundary of B2 TiNi phases, and a small amount of Ti3Ni4 phases precipitate inside B2 TiNi phases. Because of the small content of Ti2Ni phase, it is difficult to distinguish Ti2Ni from Ti3Ni4 particles in Figure 2. Therefore, in discussing Figure 2, the precipitation of Ti2Ni phase in TiNi phase is neglected. In addition, the content of Ti3Ni4 phase for the as-cast Ti49Ni51 alloy (see Figure 2a) is significantly higher than that for the heat-treated Ti49Ni51 alloy (see Figure 2b-d). In addition, the content and size of Ti3Ni4 phase decrease with the increase of heat treatment temperature, as shown in Figure 2b-d. The grain sizes of Ti3Ni4 phase for the as-cast Ti49Ni51 alloy (see Figure 2a) are also larger than those of the heat-treated Ti49Ni51 alloy (see Figure 2b-d). In fact, the range of size of Ti3Ni4 particles in the as-cast alloy is 0.2-7 µm; for the 723 K heat-treated alloys, the range of size of Ti3Ni4 particles is 0.2-3 µm; for the 773 and 823 K heat-treated alloys, the ranges of size of precipitates are 0.2-2 µm and 0.2-1 µm, respectively. Accordingly, the range of size of precipitates decrease with the increase of heat treatment temperatures. The microstructures in Figure 2 show that the grains of TiNi phase are surrounded by the continuous or discontinuous layers of Ti 3 Ni 4 phase. This phenomenon is connected with the so-called complete and incomplete wetting of grain boundaries by the second solid phase both in the Ti-Fe(Co) and Ti-Fe-Sn alloys [36][37][38]. A large and/or long grain of Ti 3 Ni 4 phase between two β-Ti phases form continuous layers, and this phenomenon is called complete grain boundary wetting (labeled in Figure 2a by the letter C); while multiple adjacent Ti 3 Ni 4 phases precipitated on grain boundaries form the discontinuous layers, corresponding to the incomplete or partial wetting of grain boundary (labeled in Figure 2a by the letter P). Therefore, in Figure 2, the Ti 3 Ni 4 phase forms the complete and incomplete grain boundary wetting.

Martensitic Transformation Temperature of Ti 49 Ni 51 Alloy
In the practical aspect of Ti-Ni SMAs, a low transformation temperature can make pipe joints, and a small thermal hysteresis can make elastic and vibrator [8]. According to the different heat-treatment conditions, Ti-Ni SMAs can undergo one-stage, two-stage, or three-stage transformation. One-stage transformation occurs mostly from the parent phase (A), with simple cubic structure (CsCl), to martensitic phase (M), with monoclinic structure (A→M); and also from A to R phase with rhomboidal structure (A→R) [39]. The two-stage transformation is from A to R phase, and then to M phase, that is, A→R→M transformation. The three-stage transformation can be divided into two cases: One is one-stage R-phase plus two-stage M transformation (R→M1→M2); the other is two-stage R phase plus one-stage M transformation (R1→R2→M) [39]. Figure 3 shows the DSC curves of as-cast and heat-treated Ti 49 Ni 51 samples. In Figure 3, M and A represent martensite transformation and reverse martensite transformation, respectively; R and R represent R-phase and reverse R-phase transformation (rhomboidal structure), respectively. The peak temperatures of transformation during cooling and heating, T M and T A , represent the martensite and reverse martensite transformation temperatures, respectively. The peak temperatures of R-phase transformation during cooling and heating, T R and T R , represent R-phase and inverse R-phase transformation temperatures, respectively. The measured T M , T A , T R , and T R values are list in Table 1. In Figure 3a, for the as-cast Ti 49 Ni 51 alloy, the transformation types are three-stage A→R, R→M 1 and R→M 2 transformation during cooling, and the two-stage A→R→M martensitic transformation occurs during heating. In fact, during the heating, the inverse transformation of the subsequent-formed martensite phase in cooling first occurs, that is, the inverse transformation of M 2 phase occurs first, compared with the M 1 phase [39]. Due to the inconsistency of the inverse transformation temperature between M 2 →R and M 1 →R, the inverse transformation peak becomes wide, or the temperature range becomes wide, as shown in Figure 3a. In Figure 3b, for the heat-treated Ti 49 Ni 51 alloy at 723 K, the two-stage A→R→M martensitic transformation occurs in the cooling stage; however, the two-stage reverse martensitic transformation is M→R→A in the heating stage. In the cooling, the T R values of the R-phase transformation temperatures are 278.8 K for the as-cast alloy, and 279.1 K for the heat-treated alloy at 723 K, and approximately the same; however, in the heating stage, the T R' values of the R-phase inverse transformation temperature is 304.9 and 285.1 K, respectively, and the T R' value of the heat-treated alloy is shifted to the low temperature, comparing with that of the as-cast alloy. The T M values of as-cast alloy and heat-treated alloy at 723 K are 211.4 and 189.1 K, respectively, which indicates that the T M of the heat-treated alloy moves to a low temperature during the cooling. But, in the heating, the reverse martensitic transformation temperature increases from 241.8 K for the as-cast alloy to 265.3 K for the heat-treated alloy at 723 K. In Figure 3c, for the heat-treated Ti 49 Ni 51 alloy at 773 K, the transformation type of the alloy during the cooling stage is the one-stage A→R transformation, while during heating, the inverse martensitic transformation is similar to the two-stage transformation of M→R→A. In Figure 3d, for the heat-treated Ti 49 Ni 51 alloy at 823 K, the phase transformation types of the alloy during cooling and heating are the one-stage A→M and M→A transformation, respectively. For the heat-treated Ti 49 Ni 51 alloy at 773 K, during the cooling, the T R value of A→R is 292.2 K, which is higher than those of the as-cast alloy and heat-treated alloy at 723 K (Seen Table 1 and Figure 2). The T R value is equal to 300 K, which is larger than that of the heat-treated alloy at 723 K, but smaller than that of the as-cast alloy.
Materials 2019, 12, x FOR PEER REVIEW 6 of 13 transformation types of the alloy during cooling and heating are the one-stage AM and MA transformation, respectively. For the heat-treated Ti49Ni51 alloy at 773 K, during the cooling, the TR value of AR is 292.2 K, which is higher than those of the as-cast alloy and heat-treated alloy at 723 K (Seen Table 1 and Figure 2). The TRvalue is equal to 300 K, which is larger than that of the heat-treated alloy at 723 K, but smaller than that of the as-cast alloy.  The AR transformation is the transition of the parent phase to a martensite phase, and it is a transition process of a simple cubic structure (BCC) with a higher symmetry to a lower symmetrical rhombohedral structure; however, the RM phase transformation is converted to a martensitic to martensitic transformation process (i.e., the symmetry of the lower rhombohedral structure to the transition of the monoclinic structure with a lowest symmetry) [31]. For the as-cast and the heat-treated samples at 723 and 773 K, the AR transformation fist occurs during the cooling. The reason is that the symmetry of the R-phase with the rhombohedral structure is higher than that of the monoclinic structure of M phase. When the temperature is further reduced, the R-phase transformation is shifted towards a monoclinic-structural martensitic phase with a lower symmetry, comparing with the symmetry of R phase (i.e., the RM transformation). For the heat-treated alloy at 823 K, the content of Ti3Ni4 phase precipitate inside and at the grain boundary of B2 phase is  The A→R transformation is the transition of the parent phase to a martensite phase, and it is a transition process of a simple cubic structure (BCC) with a higher symmetry to a lower symmetrical rhombohedral structure; however, the R→M phase transformation is converted to a martensitic to martensitic transformation process (i.e., the symmetry of the lower rhombohedral structure to the transition of the monoclinic structure with a lowest symmetry) [31]. For the as-cast and the heat-treated samples at 723 and 773 K, the A→R transformation fist occurs during the cooling. The reason is that the symmetry of the R-phase with the rhombohedral structure is higher than that of the monoclinic structure of M phase. When the temperature is further reduced, the R-phase transformation is shifted towards a monoclinic-structural martensitic phase with a lower symmetry, comparing with the symmetry of R phase (i.e., the R→M transformation). For the heat-treated alloy at 823 K, the content of Ti 3 Ni 4 phase precipitate inside and at the grain boundary of B2 phase is obviously lower than those of the as-cast alloy and the heat-treated alloy at 723 and 773 K, as shown in Figure 2. Accordingly, the microstructure of the heat-treated alloy at 823 K is in the state of recrystallizing, which results in the decrease of dislocation density, the decrease of martensitic transformation resistance [31,39], and the increase of transformation peak position. It indicates that the T M value of the heat-treated alloy at 823 K is higher than that of the heat-treated alloy at 723 K. When the heat treatment temperature is high, the density of residual defect in the alloy decreases, the microstructure uniformity improves, and the effective position of R-phase nucleation decreases, which leads to the delay of R-phase transformation or the decrease content of R-phase transformation [31,39]. The R-phase transformation is not easy to be detected on the DSC curve. Accordingly, the one-stage transformation occurs in heat-treated Ti 49 Ni 51 alloy at 823 K (i.e., the one-stage transformation of A→M and M→A).
The as-cast alloy and heat-treated alloy at 723 K exhibit high residual dislocation density, large residual stress, and more residual texture [31,39]. The interaction of these structural defects with the stress field of M transformation can inhibit the martensitic phase transformation, which results in that the martensitic transformation is delayed (i.e., the T M value becomes low) [20,26]. In addition, as the stress field generated from the R-phase transformation is weaker than that generated from the M-phase transformation, the effect of the structural defect on the stress field of the R-phase transformation is small, and the inhibition of the A to R phase transformation is reduced [31,39]. Accordingly, the R-phase transformation takes place preferentially, that is, the T R value is larger than the T M value. These will result in the separation of the R-phase and the M-phase, as shown in Figure 3a,b.
For superelastic alloys, thermal hysteresis (∆T) of phase transformation is a thermodynamic performance index in industrial applications [39]. The ∆T value is the temperature difference of the positive and negative peaks of transformation, which indicates the width of range of the operating temperature for the devices made of superelastic alloys [39]. The larger the ∆T value is, the wider the operating temperature range of the device is. The alloy with the larger ∆T value is suitable for making the connection elements, and the alloy with the smaller ∆T value is suitable for making sensors [39]. For the heat-treated alloy at 723 K, the thermal hysteresis (∆T M ) of M-phase transformation (76.2 K) is larger than that of the as-cast alloy (30.4 K), due to the decrease of the martensitic transformation temperature during the cooling, and the increase of the reverse transformation temperature during the heating for the heat-treated alloy at 723 K. For the heat-treated Ti 49 Ni 51 samples at 723 and 773 K, the thermal hysteresis (∆T R ) values of the R-phase transformation are 6.0 and 7.8 K, respectively, which are larger than that of as-cast alloy. The difference of two ∆T R values for the heat-treated samples at 723 and 773 K is small, which indicates that the thermal hysteresis of R-phase transformation has good stability at the different heat treatment temperatures. In fact, the as-cast alloy and heat-treated alloy at 723 K have the larger thermal hysteresis, which are suitable for bonding devices. For the heat-treated alloy at 773 K, only a small thermal hysteresis is suitable for sensor devices. For the 823 K heat-treated Ti 49 Ni 51 alloy, only a large thermal hysteresis is not only suitable for making junction devices, but also suitable for making sensor devices with a wide temperature range. Figure 4 shows the first cyclic loading/unloading stress-strain curve of the heat-treated Ti 49 Ni 51 alloy (723 K) at a strain of 7%. In Figure 4, the initial (σ As ) and finishing (σ Af ) stress of martensitic transformation, the initial (σ Ms ) and the finishing (σ Mf ) values of reverse martensitic transformation are labeled by the method of two tangents. The area defined by the unloading stress-strain curve and the strain is defined as the recoverable strain energy density (W r ); the area enclosed by the load stress-strain curve and the unloading stress-strain curve is defined as the dissipation energy (W d ); in the unloading process, the strain is defined as the residual strain (ε r ) when the stress is equal to zero, as shown in Figure 4.     Figure 5 shows the cyclic loading/unloading nominal stress-strain curves of as-cast and heat-treated Ti 49 Ni 51 samples at a strain of 7%. In the loading process of the as-cast Ti 49 Ni 51 alloy, the σ As values of martensitic transformation are exhibited in Figure 5a, but the σ Af values of martensitic transformation are not found in Figure 5a; during unloading, the stress platform of reverse martensitic transformation is not presented. For the heat-treated samples at 723 and 823 K, the stress transformation platform are exhibited in the loading and unloading processes, and the σ As and σ Af values can be presented in Figure 5b,d; the σ Ms and σ Mf values of reverse martensitic transformation can be presented during unloading. For the heat-treated alloy at 773 K, the stress transformation platform is obvious in the loading process, which indicates that the σ As and σ Af values can be distinguished; while in the unloading process, the strain transformation platform is not observed in Figure 5c, but the σ Ms value can be calculated.    The relation of σ As , σ Af , σ ms , σ mf values calculated from Figure 4 and cyclic loading/unloading number (n) is shown in Figure 6. In Figure 6a, the σ As values for the as-cast Ti 49 Ni 51 alloy vary greatly with the n values, while the σ As values of the heat-treated alloy change slightly. In addition, at the same loading times, the σ As values of the as-cast Ti 49 Ni 51 sample are larger than those of the heat-treated Ti 49 Ni 51 samples at three temperatures, but the σ As values of the heat-treated Ti 49 Ni 51 samples increase with the increase of the heat treatment temperatures, as shown in Figure 6a. In Figure 6b, the σ Af values of the heat-treated alloy at 723 K are approximately unchanged. Most of the σ Af values of the heat-treated alloy at 773 K are similar, except for a small amount of σ Af values. The σ Af values of the heat-treated alloy at 823 K change slightly from the third to the seventh time, while the σ Af values of the others change greatly. In fact, the σ Af values increase with the increase of the heat treatment temperatures. The relation of As, Af, ms, mf values calculated from Figure 4 and cyclic loading/unloading number (n) is shown in Figure 6. In Figure 6a, the As values for the as-cast Ti49Ni51 alloy vary greatly with the n values, while the As values of the heat-treated alloy change slightly. In addition, at the same loading times, the As values of the as-cast Ti49Ni51 sample are larger than those of the heat-treated Ti49Ni51 samples at three temperatures, but the As values of the heat-treated Ti49Ni51 samples increase with the increase of the heat treatment temperatures, as shown in Figure 6a. In . Figure 6. The relation of As, Af, ms, mf and cyclic loading/unloading number (n). (a) As and n; (b) Af and n; (c) ms and n; (d) mf and n.

Mechanical Properties of the Ti 49 Ni 51 Alloy
Due to the high density of residual dislocation, large residual stress, and large residual texture of the as-cast alloy, the resistance of M-phase transformation of the B2 TiNi structure increases under compressive stress [31], which leads to the larger initial stress and finishing stress of martensite transformation, comparing to the heat-treated alloy. Although R-phase transformation exists in as-cast alloys, it can be seen from Figure 3 that the enthalpy of R-phase transformation is obviously smaller than that of M-phase transformation. Accordingly, the M-phase transformation is the main phase transformation in the compression process, and a small amount of R-phase transformation does not significantly reduce the initial and finishing stress of martensitic transformation. For the 723 and 773 K heat-treated samples, the dislocation density, residual stress and residual texture decrease [31], and the main R-phase transformation occurs during the cooling process. When the compressive stresses are applied to the heat-treated samples at 723 and 773 K, the main R-phase transformation occurs. However, the stress field generated from the R-phase transformation is weaker than that generated from the M-phase transformation [31], which results in that the As and Af values of the heat-treated samples at 723 and 773 K are lower than the corresponding values of the as-cast sample. For the heat-treated alloy at 823 K, as the alloy is in recrystallization state, the effective position of R-phase nucleation decreases [31], and only M-phase transformation occurs, as Figure 6. The relation of σ As , σ Af , σ ms , σ mf and cyclic loading/unloading number (n). (a) σ As and n; (b) σ Af and n; (c) σ ms and n; (d) σ mf and n.
Due to the high density of residual dislocation, large residual stress, and large residual texture of the as-cast alloy, the resistance of M-phase transformation of the B2 TiNi structure increases under compressive stress [31], which leads to the larger initial stress and finishing stress of martensite transformation, comparing to the heat-treated alloy. Although R-phase transformation exists in as-cast alloys, it can be seen from Figure 3 that the enthalpy of R-phase transformation is obviously smaller than that of M-phase transformation. Accordingly, the M-phase transformation is the main phase transformation in the compression process, and a small amount of R-phase transformation does not significantly reduce the initial and finishing stress of martensitic transformation. For the 723 and 773 K heat-treated samples, the dislocation density, residual stress and residual texture decrease [31], and the main R-phase transformation occurs during the cooling process. When the compressive stresses are applied to the heat-treated samples at 723 and 773 K, the main R-phase transformation occurs. However, the stress field generated from the R-phase transformation is weaker than that generated from the M-phase transformation [31], which results in that the σ As and σ Af values of the heat-treated samples at 723 and 773 K are lower than the corresponding values of the as-cast sample. For the heat-treated alloy at 823 K, as the alloy is in recrystallization state, the effective position of R-phase nucleation decreases [31], and only M-phase transformation occurs, as shown in Figure 3d. As the stress field generated from the M-phase transformation is larger than that generated from the R-phase transformation [31], the M-phase transformation stress is higher than the R-phase transformation stress under compressive stress, which explains that the σ As and σ Af values of the heat-treated sample at 823 K are higher than the corresponding values of heat-treated samples at 723 and 773 K.
In Figure 6c, the σ Ms values of the as-cast alloy vary greatly, while the σ Ms values of the heat-treated alloy are approximately unchanged, while the σ Ms values of the as-cast alloy are higher than that of σ Ms values of the heat-treated alloy. In addition, the σ Ms values of the heat-treated alloy increase with the increase of heat treatment temperatures. In Figure 6d, the as-cast alloy only shows the three σ Mf values from first unload to third unload, but the σ Mf values of the heat-treated alloy at 773 K are not shown in Figure 5c. The variation of σ Mf values are small for the heat-treated alloy at 723 K, but the variation of σ Mf values fluctuate for the 823 K heat treatment alloy. The σ Mf values of the heat treatment alloy at 823 K are higher than those of the heat-treated alloy at 723 K, as shown in Figure 6d. In fact, the as-cast and 823 K heat-treated samples have relatively large changes in stress of martensitic transformation, and weak relative stability, which indicates that the as-cast and 823 K heat-treated samples are the hyperelastic materials with poor stability; the stress of the martensitic transformation for the heat-treated samples at 723 and 773 K changes slightly, and have relatively high stability, which indicates that the two heat-treated samples are the hyperelastic materials with good stability. Figure 7 shows the relation between the calculated σ max , ε r , W d , W r values, and cyclic loading/unloading number (n). In Figure 7a, the σ max values of the as-cast alloy increase with the increase of loading times, increasing from 781 MPa in the first cycle to 1117 MPa in the tenth cycle, and increasing approximately linearly from the sixth cycle, resulting in that the linear increase of maximum stress can easily cause plastic deformation of the sample. For the heat-treated alloy at 773 K, with the increase of cycle times, the σ max values increase rapidly at first, then slowly, and finally tend to be stable, which improves the safety during cyclic loading and unloading. Under the 723 and 823 K heat-treatment conditions, the σ max values of the alloy decrease slowly, and during the whole cycle, those values decrease slightly, and finally tend to be stable. Figure 7b shows the relation between the ε r and n values. For the as-cast alloy, the ε r values decreases rapidly from the first time to the sixth time, then tends to be stable, and are less than 0.2%. The ε r values of the heat-treated alloy at 723 K increase from 0.06% to 0.37% with the increase of cycle times. In fact, the increase is small. In addition, for the heat-treated samples at 773 and 823 K, a small ε r value appears for the first time, and the ε r values are equal to zero for the rest of the cycles. It shows that the content of martensitic transformation is approximately equal to that of reverse phase transformation under loading and unloading conditions. Figure 7c shows the relation between the W d and n values. The W d values of the as-cast sample are larger than those of the heat-treated samples, while the W d values of the 823 K heat-treated sample are significantly smaller than those of the as-cast sample and heat-treated samples at 723 and 774 K. The W d values of the 823 K heat-treated alloy begin to decrease rapidly with the increase of cycle number, and slowly decline from the fourth time; the W d value of the 723 K heat-treated alloy decrease rapidly between the first and second time, slowly decline from the second time, and finally change slightly, and tend to be stable. The W d values of the 773 K heat-treated alloy decrease slowly, and finally stabilize during the whole cycle. These results show that the W d values of the heat-treated samples at 723 and 773 K are relatively stable, though smaller than those of the as-cast sample. As the σ max values of the as-cast alloy increase gradually, the small plastic deformation occurs during the final cyclic loading process, as shown in Figure 5a. Although there are the large W d values of the as-cast alloy, it has low safety in application. As the σ max values tend to be stable for the 723 and 773 K heat-treated samples with a large W d value, the safety of the heat-treated samples is higher than that of the as-cast sample in application. Because the residual strain tends to zero in the 823 K heat-treated alloy, the smaller dissipated energy can be used in the working environment, where the dissipated energy requirement is not high.  The Wr values of the as-cast sample increase with the increase of cycle times, and are larger than those of the heat-treated samples at 723 and 773 K. For the 823 K heat-treated alloy, the Wr values decrease slowly with the increase of cycle number, and tend to remain unchanged at last. For the 723 K heat-treated alloy, the Wr values change slightly or tend to be stable during the whole cycle. For the 773 K heat-treated alloy, the Wr values change slightly from the second time, and finally tend to be stable. Therefore, the Wr values of the heat-treated samples are more stable than those of the as-cast sample. In fact, the samples treated at 773 and 823 K have relatively stable strain-energy densities, and are a good superelastic alloy.

Conclusions
The martensitic transformation and mechanical properties of the as-cast and heat-treated Ti49Ni51 samples prepared by rapidly-solidified technology were investigated. The main results are summarized as follows: (1) The microstructures of the as-cast and 723 K heat-treated Ti49Ni51 samples are composed of B2 TiNi phase, Ti3Ni4, and Ti2Ni phases; the microstructures of the heat-treated Ti49Ni51 samples at 773 and 823 K are composed of B2 TiNi phase, and of B2 TiNi phase and Ti2Ni phase, respectively. The content and size of Ti3Ni4 phase decrease with the increase of heat treatment temperatures. (2) The transformation of the as-cast Ti49Ni51 sample are third-stage ARM1 and RM2 during cooling, and two-stage ARM transformation during heating. The transformations of the heat-treated Ti49Ni51 samples at 723 and 823 K are the two-stage ARM and one stage AM transformation during cooling/heating, respectively. For the heat-treated sample at 773 K, the transformations are the AR/MRA during cooling/heating. (3) For the heat-treated alloy at 773 K, only a small thermal hysteresis is suitable for sensor devices.
For the 823 K heat-treated alloy, only a large thermal hysteresis is not only suitable for making junction devices, but also suitable for making sensor devices with a wide temperature range.  Figure 7d shows the relation between the W r values and the number of cycles. The W r values of the as-cast sample increase with the increase of cycle times, and are larger than those of the heat-treated samples at 723 and 773 K. For the 823 K heat-treated alloy, the W r values decrease slowly with the increase of cycle number, and tend to remain unchanged at last. For the 723 K heat-treated alloy, the W r values change slightly or tend to be stable during the whole cycle. For the 773 K heat-treated alloy, the W r values change slightly from the second time, and finally tend to be stable. Therefore, the W r values of the heat-treated samples are more stable than those of the as-cast sample. In fact, the samples treated at 773 and 823 K have relatively stable strain-energy densities, and are a good superelastic alloy.

Conclusions
The martensitic transformation and mechanical properties of the as-cast and heat-treated Ti 49 Ni 51 samples prepared by rapidly-solidified technology were investigated. The main results are summarized as follows: For the 823 K heat-treated alloy, only a large thermal hysteresis is not only suitable for making junction devices, but also suitable for making sensor devices with a wide temperature range.
(4) The stable σ max values of the 723 and 773 K heat-treated samples with a large W d value exhibit that the safety of the heat-treated samples is higher than that of the as-cast sample in application.
The samples treated at 773 and 823 K have relatively stable strain-energy densities, and are a good superelastic alloy.