The Investigation of Microstructure and Mechanical Behavior and the Fractographic Analysis of the Ti 49.1 Ni 50.9 Alloy in States with Different Activation Deformation Volumes

: In this article, the microstructure and mechanical behavior of the Ti 49.1 Ni 50.9 alloy with a high content of nickel in a coarse-grained state, obtained by quenching, ultraﬁne-grained (obtained through the equal-channel angular pressing (ECAP) method) and nanocrystalline (high pressure torsion (HPT) + annealing), were investigated using mechanical tensile tests at different temperatures. Mechanical tests at different strain rates for determining the parameter of strain rate sensitivity m were carried out. Analysis of m showed that with an increase in the test temperature, an increase in this parameter was observed for all studied states. In addition, this parameter was higher in the ultraﬁne-grained state than in the coarse-grained state. The activation deformation volume in the ultraﬁne-grained state was 2–3 times greater than in the coarse-grained state at similar tensile temperatures. Fractographic analysis of samples after mechanical tests was carried out. An increase in the test temperature led to a change in the nature of fracture from quasi-brittle–brittle (with small pits) at room temperature to ductile (with clear dimples) at elevated temperatures. Microstructural studies were carried out after the tensile tests at different temperatures, showing that at elevated test temperatures, the matrix was depleted in nickel with the formation of martensite twins.


Introduction
TiNi (Nitinol) alloys are an important class of shape memory alloys (SMAs). They are well known for their shape memory effect and superelasticity, and have found several important applications. These alloys are used as functional materials in mechanical actuation systems, couplings [1], actuators, biomedical equipment and biomedical implants [1,2]. In some cases, the level of properties of alloys in a coarse-grained state is insufficient. There are various methods for increasing the mechanical and functional properties of these alloys-rolling, rotational forging, and others. Methods for enhancing performance include severe plastic deformation (SPD) methods [3][4][5][6][7][8][9][10]. As a result of equal-channel angular pressing (ECAP), an ultrafine-grained (UFG) structure with a grain size of about 500 nm is formed in TiNi alloys, which leads to a significant increase in the mechanical properties and functional characteristics of shape memory effects [9,10]. Recently, many works have been devoted to the study of alloys with an ultrafinegrained structure [11][12][13]. However, the mechanical behavior and deformation mechanisms of UFG TiNi are poorly studied. Note that martensitic transformations during deformation of TiNi alloys play a strategic role in their behavior during deformation [14][15][16][17][18][19][20][21][22][23][24]. The mechanical behavior of UFG and nanocrystalline (NC) metals is markedly different from the behavior of microcrystallines (with grain sizes greater than 1 µm) [25][26][27][28][29][30][31][32]. In UFG materials, an increase in the rate of diffusion creep and an increase in the contribution of grain-boundary slip to deformation have been observed [25][26][27][28][29][30][31][32].
To assess the mechanisms of deformation, the following parameters are used-the parameter of strain rate sensitivity m and the activation volume of deformation V [24][25][26][27][28][29][30][31][32][33][34]. The value of m characterizes the ability of the material to deform uniformly, or in other words, the tendency of the material to form a neck. The physical meaning of the indicator m is that when the neck is formed, the strain rate in it will be greater than in the rest of the sample. At high values of m, this will lead to hardening of the material in the neck; therefore, the flow is concentrated in another part of the sample. Analysis of the activation volume allows us to estimate the contribution of various deformation mechanisms to a particular structural state. It is recognized that in materials with bcc and hcp lattices, the activation volume does not change with decreasing grain size, and in fcc materials, V ef decreases with decreasing grain size. As a rule, different mechanisms of plastic deformation correspond to different ranges of m, and when the mechanisms of superplastic deformation (slip due to the elongation and rotation of the grain along grain boundaries) prevail, m reaches its highest level of 0.4 -0.7 [26][27][28][29][30]. The effect superplastic deformation in alloys with a refined structure can be observed not only at elevated test temperatures. In some ultrafine-grained materials, this effect can be observed at relatively low temperatures and/or relatively high strain rate. It was shown that sensitivity to the strain rate increases with decreasing grain size for fcc metals such as Cu, Al, and Ni [30,31]. For pure Ni, for example, a decrease in grain size leads to an increase in the indicator of strain rate sensitivity m [31]. In the case of hcp metals, the dependence of m on the grain size is similar [32][33][34][35].
Experimental results on the effect of grain size and temperature on the sensitivity to the strain rate of TiNi showed significantly small effects. The mechanical behavior of TiNi alloys is studied at elevated temperatures and low strain rates that create favorable conditions for the creep of coarse-grained (CG) TiNi alloys. Previous studies have shown that the sensitivity to the strain rate of TiNi alloys strongly depends on the deformation temperature and strain rate, regardless of the alloy composition (equiatomic composition or nickel-enriched alloys) [36][37][38][39][40][41]. In materials with an increased parameter of strain rate sensitivity, the formation of a neck was not observed until fracture. Data on the change in the activation volume of the TiNi alloy, depending on the grain size, are currently not available in the literature. The austenitic intermetallide TiNi has an atomic-ordered B2 structure with a bcc type lattice; however, when cooled below M s or under the influence of stresses at temperatures below M d , the B2 phase undergoes a martensitic transformation into phase B19 with a monoclinic distorted orthorhombic unit cell.
The main goal of this paper is to elucidate the correlation between the microstructure, fracture, mechanical behavior, and the strain rate/temperature sensitivity of the TiNi alloy processed by ECAP and high pressure torsion (HPT), which have not been considered comprehensively.

Materials and Methods
Test specimens were prepared from a commercially available TiNi alloy (MATEK-SMA Ltd., Russia), manufactured using the vacuum-induction melting method. The chemical composition of the alloy was: Ti (balance); 55.08 Ni; 0.051 C; 0.03 O; 0.002 N (wt.%)resulting in the Ti 49.1 Ni 50.9 alloy (at.%) of stoichiometric composition. The temperatures of the onset and completion of the direct and reserve martensite transformation (MT) of the TiNi alloy were determined using the differential scanning calorimetry method (DSC): M S = −7 ± 2 • C; M f = −35 ± 2 • C; A s = 5 ± 2 • C; A f = 28 ± 2 • C, and this alloy has a B2 austenite structure at room temperature. The Ti 49.1 Ni 50.9 alloy was preliminary quenched from 800 • C in water. To obtain the UFG state, the Ti 49.1 Ni 50.9 alloy was subjected to equal-channel angular pressing at a temperature of 450 • C, the angle of intersection of the equipment channels was 120 • , and the number of ECAP cycles was n = 8 [9,10]. It was previously shown that an amorphous-nanocrystalline state is formed in TiNi during the high pressure torsion (HPT), and an NC state is formed during subsequent annealing [7]. Accordingly, the nanocrystalline structure was formed using the HPT method with a number of rotations of the anvils of n = 5 with a speed of 1 rpm at room temperature and a pressure of 6 GPa, and then the samples after HPT were subjected to subsequent annealing at 350 • C for 1 h.
Mechanical tensile tests under quasistatic uniaxial tension at various temperatures were carried out on a tensile testing machine designed using the Instron 5982 tensile testing machine on small flat samples with a gauge section of 1 × 0.25 × 4 mm. Tensile specimens were deformed to failure with constant crosshead speed corresponding to the initial strain rate . ε = 1 × 10 −3 at test temperatures of 25 • C, 250 • C and 400 • C. At least three samples of each type were tested. Based on the analysis of the recorded loading diagrams (stress-strain curves), in addition to the conventionally used strength and ductility properties, some functional properties of the TiNi alloy under study were determined; that is, the martensite shear stress (the critical stress to induce the B19 transformation), namely σ M . Samples for testing were cut from billets using electric-spark cutting, grinded and polished with a diamond paste. To determine the parameter of the strain rate sensitivity of the plastic flow stress m, the method of varying the tensile strain rate from . ε = 1 × 10 −3 s −1 to . ε = 1 × 10 −4 s −1 and back was applied. The parameter m was determined from Relation (1) [24]: where m is the parameter of strain rate sensitivity; σ 2 , σ 1 -flow stress, MPa;έ 2 ,έ 1 -strain rate, s −1 .
The activation volume of deformation was calculated using Formula (2) [25]: where m is the parameter of strain rate sensitivity; σ-flow stress, MPa; k-Boltzmann's constant, 1.38 × 10 −23 J/K; T-test temperature, K. The microstructure in the initial state was examined with an OLYMPUS GX51 optical microscope. The fine structure of the samples was investigated using a JEOL JEM-2100 transmission electron microscope at an accelerating voltage of 200 kV. Foils for transmission electron microscopy were obtained on a TenuPol-5 double-sided electropolishing unit using a standard electrolyte composition: 10% HClO 4 + 90% CH 3 (CH 2 ) 3 OH. To analyze the nature of the fracture of the specimens after mechanical testing, the studies were performed on a JEOL JSM 6390 scanning electron microscope (SEM) (Japan) in the secondary electron (SE) mode; the accelerating voltage was U = 10 and 20 kV.

Microstructural Studies of the Ti 49.1 Ni 50.9 Alloy
According to the data of optical metallography, the alloy has an austenitic structure at room temperature, and the grain size is 30 ± 5 µm (Figure 1). According to the TEM, the initial structure predominantly has equilibrium grain boundaries; the body of the grains is free from defects.
After ECAP, the alloy is also in the austenitic state, which was confirmed by the analysis of microdiffraction ( Figure 2). This state is characterized by a high density of dislocations in the grain body and a significant nonequilibrium character of the grain boundaries, and the average grain size after ECAP is 300 ± 40 nm ( Figure 2). A similar microstructure is formed in other TiNi alloys after ECAP [10]. The electron diffraction pattern confirms the ultrafine-grained character of the structure and the presence of only the B2 phase.  As a result of HPT at room temperature, an amorphous-nanocrystalline structure was formed in the Ti 49.1 Ni 50.9 alloy (Figure 3a), with a predominance of the amorphous phase (the volume fraction of the crystalline phase was~8 ± 2%). In the structure, the band regions with an increased fraction of the crystalline component (NC) were found to alternate with the regions of the amorphous phase (A). Probably, this structural separation is associated with the localization of deformation in shear bands. It can be seen from the microdiffraction pattern that the intense halo from the amorphous phase is imposed on weak reflections uniformly distributed over the ring from planes of the {110} B2 type. To form a fully crystalline structure, annealing was carried out at 350 • C for 1 h. As a result, the formed structure has a nanocrystalline character with an average grain size of the B2 phase of 41 ± 2 nm (Figure 3b). It is important to note that after annealing at a temperature of 350 • C, the diffuse halo is practically not visualized in electron diffraction patterns; therefore, the crystallization of the amorphous phase is completed, and the electron diffraction pattern has the form of rings with distributed point reflexes.   Analysis of the tensile curves showed that a flat area can be observed on the tensile curves of the samples tested at room temperature at stresses of σ m 350 ± 20 MPa due to the reorientation of martensite B19 under stress. When the test temperature is increased to 250 • C and higher, this flat area is absent on the tensile curves, which indicates either the blocking of the B2-B19 transformation or the B2-B19 transformation simultaneously with plastic deformation.
ECAP led to a noticeable increase in tensile strength and yield stress compared with the initial CG state, which is the result of the grinding grain. At room temperature, the tensile curves of the alloy in the CG and UFG states had an extended stage of uniform hardening, which is associated with the possibility of prolonged hardening of the martensitic phase. Tension at room temperature ended with the fracture of the samples without the formation of a neck and localization of deformation. With increasing temperature, the tensile strength and elongation to failure decreased. The HPT of the Ti 49.1 Ni 50.9 alloy led to maximum values of tensile strength of up to 1700 MPa, but at room temperature the material fractured in a brittle manner. An increase in the temperature of the test led to a decrease in strength and the appearance of noticeable plasticity. The NC alloy at 400 • C had a slightly higher ductility than CG and UFG. Perhaps this is the result of the activation of intergranular slippage. To study the characteristics of mechanical behavior, determine mechanical characteristics at elevated temperatures and determine possible deformation mechanisms, CG and UFG Ti 49.1 Ni 50.9 alloys were tested at various strain rates ( Figure 5).
The data on the speed sensitivity, presented in Table 2, show that as the temperature of the test increased, the speed sensitivity parameter increased. The parameter of speed sensitivity m in the UFG state was higher than in the CG state, and with increasing test temperature, m increased both for the CG and UFG state of the alloy. Based on parameter V, deformation occurred by the mechanism of sliding of dislocations and forest dislocations. With increasing deformation temperature, the ∆V/b 3 value increased. The increase in activation volume with increasing test temperature is probably due to the fact that at room temperature the deformation of the alloy occurs in the martensitic state, and with an increase in the deformation temperature, the martensitic transformation is blocked, and the tension goes into the austenitic state. Since the samples are fragile in the NC state at room temperature, the speed switching tests require additional research (in this regard, data on NC are not shown in the table). Since the strain rate sensitivity in the CG alloy at room temperature is low, the absence of localization and neck formation in the CG alloy in this case is apparently related to the specificity of dislocation hardening in the martensite phase. Since there is no noticeable hardening in the tensile curve of the NC alloy at 400 • C after reaching UTS, the suppressed neck formation in this case is associated with other reasons, possibly with increased strain rate sensitivity.

Fractographic Studies and Microstructure Analysis after Tensile Testing
Fractures of the specimens were examined after the mechanical tensile tests at various test temperatures (Figures 6-8). An increase in the temperature of the tests led to a change in the nature of fracture from quasi-brittle-brittle (with shallow dimples) at room temperature to viscous (with clear dimples) at elevated temperatures. This behavior is characteristic of TiNi alloys both in coarse-grained and ultrafine-grained states; however, the diameter of the dimples in the ultrafine-grained state was much smaller (the diameter of the dimples ranged from 2 to 3 microns) than in the coarse-grained (the diameter of the dimples was about 5-7 microns). In the case of a quasi-brittle and brittle fracture, the depth of the dimples was small (about 1 µm, the area on the upper right in Figure 6), and in the case of a viscous fracture (the area on the lower right in Figure 6), dimples with a depth of several µm were observed. In the GG and UFG states at elevated test temperatures, a noticeable narrowing of the sample was observed in comparison with the tests at room temperature.   In the nanocrystalline state, an increase in the test temperature, as a rule, does not lead to a change in the nature of fracture-both at room temperature and at 400 • C-the fracture pattern remains brittle with characteristic cleavage facets. Only in some areas of the fracture, dimples characteristic of viscous fracture could be observed. At the same time, it was noted that the ductility of the NC alloy at a temperature of 400 • C was noticeably higher than that of CG and UFG. Additionally, the general view of the fracture indicates that there was no noticeable localization of deformation in the NC alloy at a temperature of 400 • C (Figure 8b), and the curve indicates a longer homogeneous flow (without hardening) (Figure 4c, curve at 400 • C). By contrast, both the fractures and the shape of the curves at a temperature of 400 • C of the CG and UFG alloys indicated a strong localization of deformation and necking (Figure 4a,b, curves at 400 • C). In the case of the CG and UFG states, the curves at 400 • C after the ultimate strength showed a noticeable decrease in stress, which indicates active necking. In the case of the NC state, a noticeably less intense decrease in stress was observed, which may indicate that there is a certain mechanism, perhaps a change in the strain rate sensitivity parameter m, which actively prevents neck formation. However, in this study, it was not possible to determine this mechanism; this requires additional research.
The microstructure of the Ti 49.1 Ni 50.9 alloy after testing at room temperature and at a temperature of 400 • C in the coarse-grained state was found not to have significant differences; in some areas particles precipitated during aging were observed, which was confirmed by analysis of the electron diffraction patterns (Figure 9). In addition, there twins of martensite B19 (h M = 150 ± 10 nm) were observed (Figure 9b). In the ultrafine-grained state, the structure after testing at room temperature is typical for an alloy after ECAP with a high density of dislocations randomly distributed in the structure. After testing at a temperature of 400 • C, a clear subgrain structure appeared inside the grains obtained in the ECAP process; in addition, at this temperature, aging particles were released, and the Ni matrix was depleted, and in some areas, the martensitic nature of the structure inside the austenitic grains was observed (h M = 10 ± 5 nm), which as also confirmed by the presence of reflections in the electron diffraction pattern from phase B19 (Figure 9d).

Conclusions
In the temperature range of tensile testing from 25 to 400 • C, the values of the dislocation yield stress and the tensile strength of the UFG state of the alloy were higher than those of the CG state. The strain rate sensitivity parameter m in the UFG state was 2-2.5 times higher than in the CG state, and m also increased with increasing test temperature. The dependence of the parameter ∆V on the test temperature was similar. The activation deformation volume of the UFG state was 2-3 times higher than that of the CG state. Up to fairly high temperatures, the ultrafine-grained TiNi alloy remained strengthened; this may be related to the results in which the strain rate sensitivity indicators were close when tested at temperatures of 250 • and 400 • C. For the NC state, it is necessary to carry out additional studies, which will make it possible to reliably determine the deformation mechanism at elevated temperatures.
The performed fractographic analysis showed that an increase in the test temperature led to a change in the nature of fracture from quasi-brittle-brittle (with shallow dimples) at room temperature to viscous (with clear dimples) at elevated temperatures. This behavior is characteristic of TiNi alloys both in coarse-grained and ultrafine-grained states; however, the diameter of the dimples in the ultrafine-grained state was much smaller than in the coarse-grained state. In the nanocrystalline state, an increase in the test temperature, as a rule, did not lead to a change in the nature of fracture-both at room temperature and at 400 • C-the fracture pattern remained brittle with characteristic cleavage facets. However, in some areas of the fracture, dimples characteristic of a viscous fracture could be observed. Analysis of the structure of the TiNi samples after mechanical tests by TEM showed that at elevated test temperatures, nanoparticles of aging are precipitated and, at the same time, the matrix depletes with the formation of martensite plates.