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Materials 2019, 12(16), 2616; https://doi.org/10.3390/ma12162616

Review
Design and Development of Ti–Ni, Ni–Mn–Ga and Cu–Al–Ni-based Alloys with High and Low Temperature Shape Memory Effects
1
Laboratory of Non-ferrous Alloys, M.N. Miheev Institute of Metal Physics of Ural Branch of Russian Academy of Sciences, Ekaterinburg 620108, Russia
2
Laboratory of structural and functional steels and alloys, Institute of new materials and technologies, Ural Federal University, Yekaterinburg 620002, Russia
*
Author to whom correspondence should be addressed.
Received: 12 July 2019 / Accepted: 14 August 2019 / Published: 16 August 2019

Abstract

:
In recent years, multicomponent alloys with shape memory effects (SMEs), based on the ordered intermetallic compounds B2–TiNi, L21–Ni2MnGa, B2– and D03–Cu–Me (Me = Al, Ni, Zn), which represent a special important class of intelligent materials, have been of great interest. However, only a small number of known alloys with SMEs were found to have thermoelastic martensitic transformations (TMTs) at high temperatures. It is also found that most of the materials with TMTs and related SMEs do not have the necessary ductility and this is currently one of the main restrictions of their wide practical application. The aim of the present work is to design and develop multicomponent alloys with TMTs together with ways to improve their strength and ductile properties, using doping and advanced methods of thermal and thermomechanical treatments. The structure, phase composition, and TMTs were investigated by transmission- and scanning electron microscopy, as well as by neutron-, electron- and X-ray diffraction. Temperature measurements of the electrical resistance, magnetic susceptibility, as well as tests of the tensile mechanical properties and special characteristics of SMEs were also used. Temperature–concentration dependences for TMTs in the binary and ternary alloys of a number of quasi-binary systems were determined and discussed. It is shown that the ductility and strength of alloys required for the realization of SMEs can be achieved through optimal alloying, which excludes decomposition in the temperature range of SMEs’ usage, as well as via various treatments that ensure the formation of their fine- (FG) and ultra-fine-grained (UFG) structure.
Keywords:
thermoelastic martensitic transformation; shape memory effect; strength and ductility; TiNi; Ni2MnGa; Cu–Al–Ni alloys; structure types; parameters of microstructure; ultra-fine grain size

1. Introduction

For more than 50 years, so-called intelligent (or smart) binary and low-doped alloys based on titanium nickelide undergoing reversible thermoelastic diffusionless martensitic transformations (TMTs) attract great attention owing to the best combination of excellent structural and functional properties [1,2,3,4,5,6]. These alloys are distinguished by unusual effects of the thermomechanical shape memory (SMEs), pseudoelasticity (PE) and damping in combination with high-strength and ductile characteristics, excellent long-term strength and durability, corrosion resistance and biocompatibility. These materials are widely used in aerospace, automotive, mining, robotic, biomedical and other fields, including the creation of various micro-electro-mechanical systems (MEMS) such as sensors and actuators [2,3,4,5,6,7,8,9,10].
A unique key feature of the B2-austenite of titanium nickelide alloys in the pre-martensitic state is the total isotropic softening of all the elastic constants Cij, modules E and G, discovered and discussed first in [11,12,13,14,15,16,17,18,19,20,21] (Figure 1). In this case, according to inelastic neutron scattering there is a significant softening of the acoustic transversely polarized phonon modes TA2 <ξξ0> κ <110> e, especially, in the vicinity of the wave vectors κ at ξ = 1/3 and ξ = 1/2, which shows progress when approaching the start temperature Ms of TMTs [22,23,24,25]. Anisotropic diffuse scattering of the X-rays, electrons and the striation of tweed image contrast in the transmission electron microscopy (TEM) images are also found (Figure 2) [4,6,15,17]. These anomalous effects are caused by the lattice instability and appearance of localized quasi-elastic phonon domains that form the short order of atomic displacements (SOD). By means of the long-period modulation (LPD) in such quasi-static nanodomains, which is responsible for emerging diffraction satellites of types “1/3” and “1/2”, their structure can be reconstructed into the structure of the martensitic phases R, B19 or B19′, respectively [15,16,17,18,19]. Within the framework of this paradigm of structural-phase transformations, localized heterogeneous nanodomains with SOD and LPD in the pre-transition austenite state are considered not only as nanostructured precursors, but also as real nuclei of thermoelastic martensite crystals of the R, B19 and B19′ types at the stage preceding the TMTs themselves [4,6,17,19]. Following these views, a crystal-geometrical model of structural-phase reconstruction was proposed using the Bain deformation and shuffling shear displacements, for example, of the type {011}<011> for B2→B19 TMTs or combined displacements of the type {011}<100> + {011}<011> for B2→B19′ TMTs (Figure 3). Here, a small magnitude of the volume effect ΔV/V = (VM − VA)/VA × 100% is very important.
In fact, this model is consistent with the two-step scenario presented in [26,27,28]. In a first step, B2→B19′ reconstruction occurs through shear on the basal plane (011)[011]. In a second step, B19 transforms to B19′ martensite in a shear process on the non-basal plane (100)[011]B2. In this second process, the crystal lattice angle changes from 90 to 97 (see Figure 3c).
It should be noted that in other metastable non-ferrous alloys with qualitative agreement between the observed pre-martensitic anomalies, in contrast to titanium nickelide, the softening of the elastic constants is realized mainly by the Ziner channel C′ on the shear system {110}<110> (Figure 4a). Their changes are essentially anisotropic (A = C44/C′ >> 1), which correlates with the possibility of shuffling shear displacements of atoms and a noticeable softening of the phonon modes of the TA2 branch [19,29,30,31,32]. For example, for the copper-based alloys we have A = 12 − 13, and for the ferromagnetic intermetallic compound Ni2MnGa, A = 23 [33,34]. It is important that in the Ti–Ni–Cu-based alloys the parameter of anisotropy increases also almost twice in comparison with that of a binary TiNi compound (Figure 4b).
In recent years, a number of papers have widely discussed the important role of the weighted average number of valence electrons per atom ev/a and their concentration cv = ev/et (where et = Z is the total number of electrons) in the temperature–concentration behavior of temperatures of TMTs as the factors responsible not only for elastic properties, but also for the overall stability of the atomic crystal lattice of the austenite in relation to the TMTs [35,36,37,38]. It was suggested that variation in the electronic band structure must be explicitly considered [38]. Valence electrons govern bond formation and thus the elastic properties [35,36]. In particular, on a large number of different doped alloys (more than 200 different alloys with TMTs) it was found that their martensite start temperatures Ms decrease almost linearly with increasing cv, although the data show large scatter for different alloys [37,38]. Conceptionally, cv mixes parameters such as the valence electron number and atomic size [38]. The martensite temperatures Ms, Mf, As, Af and their dependence on alloy composition are of the utmost importance in shape memory technology [1,2,3,4,5,6,7,8,9,10,21,28,32,38].
As is known, TMTs and related SMEs in the binary and low-doped virtually equiatomic Ti–Ni alloys are realized in the operating temperatures below 373 K, limiting or excluding some useful practical applications of them. First, this is due to the “wedging” of the homogeneity region of existence of the phase B2 in the phase equilibrium diagram to the stoichiometric composition of Ti50Ni50 below 900 K, which leads to decomposition of non-stoichiometric Ti–Ni-alloys during heat treatment.
Second, the changes of the chemical composition of binary Ti–Ni alloys with the SMEs are limited between 50.0 and 51.2 at.% Ni due to the strong concentration dependence of the start (Ms, As) and finish (Mf, Af) temperatures of the forward and reverse TMTs. They are almost linearly reduced by more than 100 K with gradual addition to 1 at.% Ni [1,2,3,4,5,6,32,38]. It is also well known that small impurities of O and C lead to a decrease in the TMTs temperatures in TiNi-based alloys and worsen their physical and mechanical properties because of forming Ti4Ni2Ox oxides and TiC carbides [1,2,3,4,5,6].
The strong influence of the concentration fluctuation of the basic components on the properties is not typical of commercial structural and functional materials. On the contrary, these materials must have a weaker concentration dependence of their exploitation characteristics and a relatively large composition tolerance, for example, in steels of no less than of 0.5–1.0 at.% content variations, which, as a rule, does not require constant precision chemical control and is technologically quite simply achieved [38,39]. Therefore, at least in the last two decades, active investigations have been conducted to find ternary and multicomponent compositions of alloys with SMEs as candidates capable of expanding and surpassing TiNi-based alloys’ unique capabilities. Doping by the various chemical elements results in different influence on the TMTs and SMEs in B2–TiNi alloys. Their solubility varies in wide limits and depends on the type of replacement Ti- or Ni-sublattice in the initial B2 superstructure into which the substitution is occurred in the binary equiatomic alloy TiNi—by the sections Ni50Ti50−xMex, Ti50Ni50−yMey or by the other way [4,6,21,31,32,38,40,41]. The analogous data were received for ordered Cu–Al–Ni and other Cu-based alloys [1,2,3,4,5,6].
Ternary alloys of Heusler, of L21 type, based on Ni2MnGa represent another important class of low-modulus ferromagnetic alloys with TMTs, in which the SMEs and PE can be induced by a magnetic field along with thermal and mechanical action [33,34,42]. For these brittle alloys, the problem of complex doping and optimization of their microstructure to provide the necessary TMT temperatures and the required physical and mechanical properties is even more necessary.
The behavior of key SME characteristics of alloys with concentration deviation from the quasibinarity is practice-important. In commercial synthesis of multicomponent materials it is difficult to obtain large-volume ingots with high-precision chemical composition and at the same time to avoid its uncontrolled deviation. Of course, it is necessary to take into account this fact in the design and development of new alloys, for example, because of a possible decomposition that determines significant changes in their physical and mechanical properties and, especially, ductility, in nonstoichiometric TiNi-based alloys [32,38,42,43,44,45,46]. In the alloys with high-temperature SMEs, the effect of homogeneous and/or heterogeneous decomposition is also responsible for changes in the properties during their exploitation [47,48,49,50]. Moreover, along with the doping and homogenizing treatment of the initial alloys, for the effective controlling of the properties one usually tends to employ—in the capacity of the economic alternative—the modern progressive methods that could radically change the alloys’ microstructural state (grain size, substructure and inclusions of excess phases). At the same time, on the one hand, various mechanisms of hardening (solid solution, deformation, aging or grain refinement) can be additionally implemented, and, on the other hand, some negative effects of alloys in the initial cast state can be weakened or modified.
The aim of the present work is the comparative systematic study of the structural and phase transformations, physical and mechanical properties of binary and ternary alloys of quasi-binary compositions with SMEs, based on the B2–TiNi, L21–Ni2MnGa and D03–CuAlNi compound systems, depending on high-doping and employment of extreme thermal and mechanical treatments with the use of rapid quenching from the melt (RQM) or severe plastic deformation (SPD).

2. Materials and Methods

The initial components of high-purity better than 99.95 wt.% were chosen for synthesis of the alloys. The alloys were obtained by the arc melting method in a helium atmosphere. Table 1 and Table 2 represent the studied binary and ternary alloys based on titanium nickelide, and Table 3 and Table 4 ternary alloys based on Ni2MnGa and Cu–Al–Ni (concentrations are given in at.%). All the alloys were re-melted three times to ensure a uniform chemical composition of the ingots. In our homogenization procedure we used hot pressing or forging of ingots by 5–10% at 1173–1273 K, subsequent annealing in vacuumed quartz capsules at 1223 K for 24 h, and quenching in water. After this treatment, generally homogeneous fine-grained (FG) microstructures (with grain sizes less than 100 µm) were obtained in the Ti–Ni-based alloys, in distinction from large-grained (LG) brittle homogenized as-cast Ni–Mn–Ga-based alloys. SPD was provided by equal-channel angular pressing (ECAP), high-pressure torsion (HPT) and multi-pass extrusion into a wire or strip. RQM ribbons with a thickness of 30 µm and a width of 1.5 mm were manufactured by quenching from the melt on a rotating copper disc, providing the cooling rate of 105–106 K/s. The structural-phase composition was determined by X-ray diffraction (XRD) in monochromatic radiation CuKα. Micro-structural and micro-chemical analysis was performed using the transmission and scanning electron microscopes (TEM) Tecnai G2 30 and (SEM) Quanta 200 Pegasus equipped with an analytical energy dispersive X-ray spectrometry (EDS) system. TEM studies included direct observation and in-situ experiments on heating or cooling. In SEM studies, EBSD and fractography of samples subjected to tensile testing were also used. TMTs temperatures (Ms, Mf, As and Af) were measured by temperature dependences of the electrical resistance and magnetic susceptibility. Mechanical properties (σm, σy, σu, σr, δ, ε and Ψ) were determined on Instron-type tensile devices at room temperature. SMEs were fixed by heating the samples to temperatures above Af, restoring their shape after bending by means of special device in the martensitic state.

3. Results and Discussion

3.1. Temperature and Concentration Dependences of TMT Temperatures

The obvious necessary conditions for the synthesis of the alloys with SMEs based on titanium nickelide are the preservation of the solid solution state in the high-doped replacement-undergoing sublattice and in general in the B2 superstructure, as well as its ability to TMTs. The ternary alloys selected for systematic study, the matrix of which was as close as possible to quasi-binary systems in chemical composition (taking into account the possible precipitation of the TiC and Ti4Ni2Ox phases, increasing in general by 0.2–0.5 at.% concentration of nickel in the austenitic B2 matrix) corresponded to such conditions. Table 1 and Table 2 show the TMTs temperatures Ms, Mf, As and Af of the binary and ternary quasi-binary alloys of titanium nickelide considered in the present paper (concentrations are given in at.%), average numbers of valence electrons per atom (ev/a) and their concentration cv = ev/z. The alloys for the study were selected so that the nominal content of their B2 austenite did not differ from the real one by more than ±0.1 at.%. From the Table 1 and Table 2 it followed that in dependence of the doping of the third component the quasi-binary alloys of titanium nickelide underwent TMTs and, respectively, SMEs in a wide range of temperatures, from cryogenic to 1000 K, and all these determined their unique practical potential.
According to XRD, TEM and SEM, it was found that all the alloys were in a single-phase austenite state (in the presence of a certain small amount of carbides and titanium oxides, undetectable by XRD). Thus, high-doped alloys of studied compositions were solid solutions and it could be assumed that the filling of the alloying addition occurred purposefully in one of replacement-experiencing Ti or Ni sublattices of the B2 superstructure TiNi. In accordance with data from Table 1 and Table 2 we replaced Ni by the Al, Mn, Fe, Co, Cu, Pd, Pt and Au or, on the contrary, Ti by the Al, Mn, Zr and Hf. The results show a pronounced composition dependence of not only TMTs temperatures, but also the number of the valence electrons and their concentration. Adding chemical elements that in the periodic table are right below Ni (period VIIIB: Pd or Pt) or Ti (period IVB: Zr and Hf) resulted in the strong increase of Ms, Mf, Af and As. On the contrary, the addition of chemical elements that in the periodical table are located between Ti and Ni (Cr [38], V [38], Mn, Fe, Co) resulted in a decrease of TMTs temperatures.
The presence of a two-stage TMT B2 ↔ R ↔ B19′ with a sharp decrease of the temperatures for the last transformation R ↔ B19′ was common to the alloys, and is given in one Table 1. The coincidence of the trends of their changes with changes of ev/a and cv, as well as in the high-doped alloys in a solid solution state on the base of L21–Ni–Mn–Ga and D03–Cu–Al–Ni (Table 3 and Table 4, Figure 5), was unusual and contradicting to the conclusions of the valence electron concentration hypothesis proposed by Zarinejad and Liu [35,36,37]. The only correlation in the changes of the martensite temperatures and cv in the TiNi-based alloys doped by Mn, Cr and V (replacing Ti) was in accordance with the data [37] (see Table 1 and Table 5). We used data from [38] on high-purity titanium nickelide alloys doped with Cr, V or Cu, and from [43] on conventional alloys, with Cu for comparison.
The behavior of the TiNi alloys doped with Cu, Pd, Pt and Au (replacing Ni) or Zr and Hf (replacing Ti; see Table 2) was even more unusual. In the case of Cu-, Zr- and Hf-doped alloys undergoing TMTs B2 ↔ B19′ the opposite changes of martensite temperatures and electron concentrations did occur. However, in the alloys moderate-doped with Pd, Pt and Au the first B2 ↔ B19′ TMTs were distinguished by coincidence of its changes. The opposite changes of martensite temperatures occurred during the second TMTs B2 ↔ B19. Adding Cu up to 15 at.% decreased all martensite temperatures of TMTs B2 ↔ B19′, but slightly increased the temperatures of the TMTs B2↔B19 (Table 2). Starting from Cu 25 at.%, the Ti50Ni50−xCux alloys undergo eutectoid decomposition with the formation of equiaxial and lamellar dispersed particles of the B11-TiCu phase [44,45].
It should also be noted that when doping with the third component, it is performed from common Ti and Ni periods IVB and VIIIB of the periodic table, that is, ev/a = 7, the TMTs temperatures increased as much as possible. However, cv from the value 0.28 corresponding to Ti50Ni50 could decrease in different ways correlated with the position of the element in the large periods of the periodic table and, consequently, with the number Z of the chemical element. Discussing the quantitative dependence of the martensite temperatures, ev/a and cv, we could still point to their significant numerical differences for various alloys with similar temperatures of TMTs. For example, for the B2–NiMnGa-based and D03–CuAlNi-based alloys they were much larger than for alloys based on the B2-compound TiNi (Table 3 and Table 4).
The influence of the alloy decomposition at temperatures of TMTs is a known important fact noted already in the discussion on the B2 alloys Ti–Ni–Cu [44,45]. At the same time, it is obvious that not only the mechanical properties of the alloys were changed, but also the temperature characteristics of the TMTs decreased. This is indicated by the data of Table 2 on the change in the temperature–concentration dependence of TMTs at a concentration of Cu more than 25 at.%, equal to the solubility limit of Cu in B2–TiNi. Such behavior of alloys was also found in varying the content of not only Cu, but also Ti and Ni in the alloys of this system, which led to decomposition [44,45]. Significant influence of small impurities O and C on the TMTs, structure and properties of titanium nickelide binary alloys was already noted. Their great role in Ti–Ni–Cu alloys with TMTs was evidenced from the data of Table 5. Thus, alloys of Ti–Ni–Cu of the same chemical composition deteriorated more by C and O having lower TMT temperatures, which could be explained only by their solid solution effect. The negative role of inclusions of oxides and carbides was great in the ductile behavior of alloys, increasing their brittleness.

3.2. Crystal Structure and Morphological Regularities of Alloys with TMTs

Table 6 and Table 7 show the effect of doping on the phase composition, on the lattice parameters of the austenitic and martensitic phases and on the volume TMTs effect, ΔV/V, with comparison to the known literature data [46,47,48]. The given results demonstrate regular concentration dependences of the martensite structural type, parameters of their unit cell and rather small values |V/V| ≤ 1%, typical of TMTs.
According to SEM and TEM investigations, the packet morphology of pairwise-twinned martensitic crystals with coherent boundaries was typical of the martensitic microstructure of all studied quasi-binary fine-grained (FG) alloys (Figure 6). As is seen from Figure 6b,c, the pressure-induced martensitic crystals in low-module alloys would be oriented in each grain in the preferred crystallographic directions, demonstrating the microstructural mechanism of the pseudoelasticity effect due to macro deformation.
The flat coherent boundaries of crystals of B19′ martensite of different orientations within the packages occur, as a rule, along the planes (011), as well as (111) and (113)B19′ [40,41,49,50,51,52,53,54]. The packet character of the morphology of the B19′ crystals pairwise-twinned by I-type (011)B19′ was clearly visible on TEM images at such grain orientations of twinned B19’ crystals when the zone axes (z.a.) of reflecting planes was [100]B19′ (Figure 7a and Figure 8a). Note also that thin secondary nanotwins were often clearly visualized within the wider initial twinned crystals of the packets.
Figure 7c, Figure 8b,c and Figure 9 present another typical example of TEM images and SAEDs of twinned martensite of packet morphology of twinned martensite, when zone axes were [110]B19′. In this case, parallel-plate or wedge-shaped crystals of the initial orientation forming the package took turns with crystals of twins of I-type (011)B19′ and II-type <011>B19′. There were martensite packets with thin-twinned crystals or with wider plates, internally thin-twinned already on secondary systems of twinning, oriented at an angle to the boundaries of primary crystals (Figure 7c and Figure 8a,c). There were nanotwins of I-type on (111) (Figure 9) and on (001) B19′ (Figure 7c, Figure 8b,d and Figure 9b). In this case, in SAED (Figure 7d, Figure 8d and Figure 9d) sharp streaks correspondent to them in the directions of the reciprocal lattice [111]*B19′ and [001]*B19′, respectively, were observed. It was found that the orientation ratio (o.r.) of B2 austenite and B19′martensite crystals was close to Bain one:
(100)B2||(100)B19′; [011]B2||[010]B19′; [011]B2||[001]B19′
As is known, in the B19′ martensite substructure of both binary and ternary titanium nickelide alloys, the twin mode of II type <011> plays an important role, responsible for the homogeneous shear deformation necessary to retain the macroscopic invariance of the habit plane at the TMT B2 → B19′ in accordance with the concepts of phenomenological crystallographic theory [47,49,51,52]. In B19′ martensite alloys with Cu, Zr, Hf, along with twins of II-type <011>B19′ a large number of twins of I-type (011) and (111)B19′, as well as composite twins on (001)B19′ were present. Obviously, the presence of all of these mentioned nanotwins of the I and II types provided both geometrically necessary shear with invariant lattice in the TMTs B2 → B19′ and pseudo-elastic volume and planar accommodation of the martensitic crystals due to the packet pair-twinned morphology and, if necessary, secondary nanotwinning in these low-modulus alloys of titanium nickelide [49,51,52]. This can explain the variety of observed variants of coherent boundaries and twinning of martensite crystals. It is assumed that pseudo-elastic composite nanotwins on (001)B19′ are often observed both in the depleted and enriched TiNi-based alloys, which are not geometrically necessary, and also provide for additional elastic-plastic accommodation of the lattice of B19′ martensite [40,41,48,49].
B19 martensite formed in titanium nickelide alloys doped with Cu, Pd, Pt and Au also had a predominantly packet-pyramidal morphology of pairwise-twinned plate crystals (Figure 10a–d). Secondary nanotwins appeared under subsequent cooling in B19 crystals or at B19 → B19′ TMTs (Figure 10e,f). Similar results were obtained on B19 martensite in alloys doped with Pd, Pt and Au when the temperature was reduced down to room temperature. The orientation ratio of the lattices of the B2 austenite and the B19 martensite was similar to the o.r. (1).
A fine structure and morphology of the packet 2 M and long-period 10 M and 14 M martensites in large-grained (LG), FG, and ultra-fine-grained (UFG) alloys based on Ni2MnGa (Figure 11 and Figure 12) and 2 H and 18 R long-period of martensites in alloys based on Cu–Al–Ni were qualitatively similar [41].
An important feature of titanium nickelide doped by copper instead of nickel in concentrations exceeding 23–25 at.%, hafnium and zirconium instead of titanium exceeding 10–12 at.% is their ability to synthesize the alloys in the amorphous state by spinning from the melt [55,56,57,58,59,60,61,62,63,64,65,66,67,68,69]. As a result of the subsequent optimal heat treatment, it was possible to provide the formation of predominantly single-packet martensite with high-temperature SMEs, a high-strength and ductile UFG structure (Figure 13).
Another effective external influence for the radical grain-size refinement of up to amorphization is provided by various methods of plastic deformation including HPT, ECAP, multi-pass cold rolling, cold drawing, local canning rolling or compression, plain strain compression, screw extrusion, and others [70,71,72,73,74,75,76,77,78,79,80,81,82,83,84,85,86,87,88,89,90,91]. Figure 14 illustrates the main types of microstructures obtained in the HPT and ECAP.

3.3. Mechanical Properties of Fine- and Ultrafine-Grained Alloys

Table 4 presents the tensile mechanical properties of the quenched Cu–Al–Ni-based alloys with different grain sizes, Table 8—FG titanium nickelide alloys of normal purity of O and C, including the dopants Cu, Fe, Co and Pd, at room temperature, and Table 9 for comparison—binary FG alloys of high purity in C and O. At close values of the average grain size (50–70 µm) the alloy Ti49.4Ni50.6 had almost twice the larger relative elongation (δ = 75%), than the best in Table 8 in the properties of the Ti50Ni50 alloy (δ = 40%). It is seen that the doping of Cu, Fe, Co and Pd of alloys led to a decrease in the relative elongation, although their strength characteristics and SMEs remained attractive. The strongest concentration changes in the mechanical properties and grain size, accompanied by sharp embrittlement, were demonstrated by copper alloys (Table 4).
A noticeable improvement in the strength and ductile properties of titanium nickelide was achieved in the formation of the UFG structure that was formed via in advanced deformation thermal technologies using SPD (ECAP, HPT, multi-pass rolling and drawing into a strip, rod or wire) of initial large-size samples of Ti–Ni alloys (Table 9) [51,52,53,54,55,56,57,58,59,60,61,62].
RQM methods can be most useful for creating thin UFG tapes with SMEs as sensors or actuators in miniature MEMS devices [55]. Their usage after additional subsequent annealing allows obtaining long homogeneous tapes with the UFG structure, attractive in mechanical properties and parameters of the SMEs. A complex of high mechanical properties of the rapidly quenched UFG alloy doped with copper (σu, 850–1550 MPa; σy, 620–1200 MPa; σm, 100–50 MPa; δ, 9–12%; reactive stress σr = σy−σm, 620–1110 MPa and a reversible deformation of the ε, 3–5%; Table 10) was received and a new method of obtaining high strength UFG SMEs alloys in the form of thin ribbons was proposed, based on the technology of spinning from the melt of the non-stoichiometric alloys Ti50+xNi25−xCu25 (x ≤ ±1 at.%) and Ti50+yNi25Cu25−y (y ≤ ±1 at.%) [44,45]. Similar structural changes needed for increasing strength and ductile properties were also found on Ni–Ti–Zr, Ni–Ti–Hf and (Ni, Cu)50(Ti, Hf)50 alloys [56,57,58,59,60,61,62,63,64]. Thus, it was found that the creation of the FG and UFG structure could significantly improve or (in the case of strong strengthening due to extreme external influences) preserve the ductility of the alloys necessary for the implementation of SMEs.
In conclusion, the results of fractographic SEM studies of alloys with TMTs were studied. It was found that in the quenched FG alloys Ti50Ni50 and Ti49.4Ni50.6 the fracture had generally a viscous transgranular character (Figure 15a). The creation of the UFG structure did not change the type of fracture and the character of the destruction of these alloys. Many centers of localization of deformation with the appearance of small flat pits and, consequently, low ridges of separation, which was typical of the viscous mechanism of fracture with low energy, were observed on the surface of the fractures (for example, Figure 15b). The quenched large-grained (LG) TiNi-based alloys doped with Cu, Zr and Hf underwent the quasi-brittle fracture. There took place classic brittle fracture with predominant plane cleavage along the grain boundaries in the Cu–Al–Ni-based alloys (for example, Figure 15c).
In our opinion, the increase of the elastic anisotropy in the high-doped alloys based on TiNi was the main reason for the strong dependence of their mechanical properties on the grain size. All other alloys with TMTs of Ni–Mn–Ga and Cu–Al–Ni systems also were characterized by the utmost high elastic anisotropy. It was because of the coherent accommodation of elastic stresses induced by the volume effect; they were localized at the grain boundaries, which ultimately becomes the general cause of intergranular brittleness of the alloys with large grain size, undergoing TMTs.

4. Summary and Conclusions

In the present work, ternary close to quasi-binary alloys based on the Ti–Ni–X system (where X = Al, V, Cr, Mn, Fe, Co, Cu, Zr, Pd, Hf, Pt or Au) and, for comparison, alloys of two actual systems Ni–Mn–Ga and Cu–Al–Ni were selected to investigate the possibilities for the creation of strength and ductile materials with TMTs and related SMEs due to multicomponent alloying and extreme external treatments. TMT temperatures, the type and parameters of the martensitic phases, the volume effect of the phase transition, the average number of valence electrons (ev/a) and their concentration (cv) were determined. A detailed study of the fine structure and morphological features of martensite in large, fine and ultrafine-grained alloys (after their synthesis by rapid quenching from the melt or severe plastic deformation) was carried out. Tensile mechanical properties of a number of alloys were measured and mechanisms of their destruction were established. From the analysis of the obtained results, taking into account the literature data, the following conclusions could be made:
  • The small volume effect |ΔV/V| < 1%, typical of TMTs, as well as the presence of pre-martensitic softening of elastic constants and the formation of a special heterogeneous pre-martensitic state was common for all studied alloys with a sufficiently noticeable temperature–concentration changes of crystal-structure parameters of austenite and martensitic phases, structural types of martensite and their metastable long-period variants.
  • All studied alloys with TMTs combined the similarity of the pair-twinned packet morphology and microstructural hierarchy, proximity to Bain ratio, the action of several geometrically necessary twinning systems of martensite and the presence of coherent variants of their boundaries and subboundaries. In the condition of low elastic constants this diversity of substructure elements in thermoelastic martensite provided pseudo-elastic planar and volume lattice accommodation, and in the case of external mechanical influences, their mobile favorable reorientation and the effect of pseudoelasticity.
  • The preservation of their single-phase high temperature austenitic state capable of TMTs played an important role in the multicomponent doping of these alloys. The chemical elements Zr and Hf from the Ti-period IVB of the periodic table had increased solubility (up to 25–30 at.%) replacing Ti in the quasi-binary alloys NiTi–NiMe. The elements Fe, Co, Pd and Pt from the Ni-period VIIIB had unlimited solubility replacing Ni in the quasi-binary alloys TiNi–TiMe. The solubility of the remained metals in the B2–TiNi compound was relatively low (up to 10 at.%), except for the unlimited solubility of Au and high (up to 25 at.%) one of Cu. It is essential that the high doping by Zr, Pd, Hf, Pt and Au provided a noticeable increase in the temperatures of TMTs, extending the temperature range of the implementation of TMTs above 373 K up to 1000 K. While Fe, Co and a number of other transition metals V, Cr and Mn, located between the elements Ti and Ni, caused them to noticeably decline below room temperature.
  • When comparing the concentration changes of TMTs temperatures, average concentrations of valence electrons (cv) and the number of valence electrons (ev/a) for the alloys of the studied doping systems, a wide variety of them were established, which in general did not allow us to correctly determine their correlation. For example, for a number of alloys with high-temperature TMTs high-doped with Zr and Hf (replacing Ti) or Cu, Pd, Pt and Au (replacing Ni), the strong increase in TMTs temperatures was indeed consistent with the opposite decrease in the cv value. There was also the divergence between these factors and doping with V, Cr and Mn. On the contrary, there was a coinciding change in the other TiNi-based alloys moderately doped by the elements Al, Fe, Co, Cu, Pd, Pt and Au replacing Ni. In L21–Ni–Mn–Ga and D03–Cu–Al–Ni alloys, a strong increase in TMTs temperatures was also accompanied by a noticeable increase in cv and ev/a.
  • The strong dependence on the grain size characterized the mechanical behavior of the studied alloys during tensile tests, especially for the relative elongation. The strength and the ductile properties of the ternary alloys decreased and the fracture became more brittle intercrystalline, in contrast to the viscous transcrystalline character in the binary TiNi alloy.
  • In our opinion, the increase of the elastic anisotropy in high-doped alloys based on TiNi was the main reason for the strong dependence of their mechanical properties on the grain size. All other alloys with TMTs of Ni–Mn–Ga and Cu–Al–Ni systems also were characterized by the utmost high elastic anisotropy. It is because of the coherent accommodation of elastic stresses induced by the volume effect, they were localized at the grain boundaries, which ultimately becomes the general cause of intergranular brittleness of the alloys with large grain size undergoing TMTs.
  • The strength and ductility of alloys with TMTs required for the realization of SMEs could be achieved by doping, which would provide a decrease in the elastic anisotropy parameter, and on the other hand, due to the formation of a fine-grained structure. Important especially for alloys with high temperature SMEs was the presence of homogeneous and heterogeneous decomposition with a strong strengthening effect that one should avoid using the quasi-binary precision doping.
In conclusion, based on the obtained experimental results we suggest dividing various shape memory alloys to two groups. The first group shows an extremely strong dependence of martensite temperatures on the alloy composition. This is, for example, observed for high-doped ternary Ti–Ni–Pd, Ti–Ni–Pt, Ti–Ni–Au, Ni–Ti–Zr and Ni–Ti–Hf alloys with high-temperature TMTs. The behavior of these alloys was well predicted by the valence electron concept proposed by Zarinejad and Liu [35,36,37] and a gradual change in the enthalpy of transformation [38]. In these alloys the B2 phase cannot be stabilized by the formation of antisite defects, because the alloy additions also stabilize martensitic phases B19 or B19′ [38]. On the contrary, the second group of Ni–Ti, Ni–Mn–Ga and Cu-based intermetallic alloys showed opposite dependence of martensite temperatures on the alloy composition and electron concentration when the trends of their changes coincided (decreased or increased together). This was observed for the binary TiNi and low and moderately doped ternary Ti–Ni–Al, Ti–Ni–Fe, Ti–Ni–Co, Ti–Ni–Cu, Ti–Ni–Pd, Ti–Ni–Pt and Ti–Ni–Au. The compositional dependence of martensite temperatures in these alloys can be rationalized on the basis of a strong stabilization of B2-austenite through the formation of antisite defects [38].

Author Contributions

Conceptualization, methodology, analysis, writing—review V.P.; methodology, investigation, analysis, writing—original draft preparation N.K., E.M. and A.P.

Funding

This work was performed within the framework of state task “Structure”, grant no. AAAA-A18-118020190116-6 and the cooperative laboratory of the Ural Federal University n.a. the First President of Russia B.N. Yeltsin and the Mikheev Institute of Metal Physics of the Ural Branch of the Russian Academy of Sciences.

Conflicts of Interest

The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

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Figure 1. Temperature–concentration dependences of elastic constants: (a) C11, (b) C′, (c) C44 and (d) elastic anisotropy parameter A = C44/C′ of the single-crystalline alloys: 1—Ti50Fe50, 2—Ti50Ni25Fe25, 3—Ti50Ni35Fe15, 4—Ti50Ni40Fe10, 5—Ti50Ni45Fe5, 6—Ti50Ni48Fe2, 7—Ti50Ni50 and 8—Ti49Ni51 [11]. (Reproduced with permission from Khachin V.N.; Muslov S.A.; Pushin V.G.; Chumlyakov Y.I., 1987, DAN SSSR).
Figure 1. Temperature–concentration dependences of elastic constants: (a) C11, (b) C′, (c) C44 and (d) elastic anisotropy parameter A = C44/C′ of the single-crystalline alloys: 1—Ti50Fe50, 2—Ti50Ni25Fe25, 3—Ti50Ni35Fe15, 4—Ti50Ni40Fe10, 5—Ti50Ni45Fe5, 6—Ti50Ni48Fe2, 7—Ti50Ni50 and 8—Ti49Ni51 [11]. (Reproduced with permission from Khachin V.N.; Muslov S.A.; Pushin V.G.; Chumlyakov Y.I., 1987, DAN SSSR).
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Figure 2. Typical image of (a) tweed TEM contrast and (bd) diffuse scattering of single-crystalline alloys of titanium nickelide; (b) X-rays and (c,d) selected area electron diffraction (SAED) [4,6].
Figure 2. Typical image of (a) tweed TEM contrast and (bd) diffuse scattering of single-crystalline alloys of titanium nickelide; (b) X-rays and (c,d) selected area electron diffraction (SAED) [4,6].
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Figure 3. Unit lattice cells of the (a) B2-, (b,c) B19- and (d) B19′ phases in titanium nickelide alloys, their dimension-orientational ratios and schemes of rearrangements determined by shuffling (of type {011}<100> and {011}<011>) atomic displacements (shear basal planes {011}B2 shaded) [17] (Copyright, 1989, Physics of Metals and Metallography).
Figure 3. Unit lattice cells of the (a) B2-, (b,c) B19- and (d) B19′ phases in titanium nickelide alloys, their dimension-orientational ratios and schemes of rearrangements determined by shuffling (of type {011}<100> and {011}<011>) atomic displacements (shear basal planes {011}B2 shaded) [17] (Copyright, 1989, Physics of Metals and Metallography).
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Figure 4. Temperature–concentration dependences of single-crystalline elastic constants C11 (dot line), C′ (dotted line), C44 (dash-dotted line) and the parameter of the elastic anisotropy A = C44/C′ (solid line); (a) BCC alloys: 1—Au-Zn-Cu, 2—Cu-Zn, 3—NiAl, 4—Cu-Al-Zn, 5—Au-50 at. % Cd and 6—Au-47.5 at. % Cd [19] and (b) 7—Ti49Ni51, 8—Ti50Ni48Fe2 and 9—Ti50Ni38Cu10Fe2 [12,13] (Copyright, 1987, Izvestia VUZov).
Figure 4. Temperature–concentration dependences of single-crystalline elastic constants C11 (dot line), C′ (dotted line), C44 (dash-dotted line) and the parameter of the elastic anisotropy A = C44/C′ (solid line); (a) BCC alloys: 1—Au-Zn-Cu, 2—Cu-Zn, 3—NiAl, 4—Cu-Al-Zn, 5—Au-50 at. % Cd and 6—Au-47.5 at. % Cd [19] and (b) 7—Ti49Ni51, 8—Ti50Ni48Fe2 and 9—Ti50Ni38Cu10Fe2 [12,13] (Copyright, 1987, Izvestia VUZov).
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Figure 5. The temperature-concentration diagram of magnetic and martensitic transformations of Ni50+xMnxGa25 (left of x = 25) and Ni50MnxGa50−x-based alloys (right). The phase fields of paramagnetic (P) and ferromagnetic (F) austenite (A) and martensite (M) are indicated.
Figure 5. The temperature-concentration diagram of magnetic and martensitic transformations of Ni50+xMnxGa25 (left of x = 25) and Ni50MnxGa50−x-based alloys (right). The phase fields of paramagnetic (P) and ferromagnetic (F) austenite (A) and martensite (M) are indicated.
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Figure 6. SEM images of the microstructure of Ti49.4Ni50.6 alloy: (a) In the austenitic state and (b,c) in the single-packet martensite state after treatment with a pressure of 6 GPa.
Figure 6. SEM images of the microstructure of Ti49.4Ni50.6 alloy: (a) In the austenitic state and (b,c) in the single-packet martensite state after treatment with a pressure of 6 GPa.
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Figure 7. TEM images of the microstructure of B19′ martensite and the corresponding SAED (a,b) of alloys Ti50Ni45Cu5 at 250 K and (c,d) Ti50Ni40Cu10 at 120 K.
Figure 7. TEM images of the microstructure of B19′ martensite and the corresponding SAED (a,b) of alloys Ti50Ni45Cu5 at 250 K and (c,d) Ti50Ni40Cu10 at 120 K.
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Figure 8. TEM images of the microstructure and the corresponding SAED of alloys Ni49.5Ti50.5−xZrx: (a) x = 3 at.%, (b) x = 10 at.% and (c,d) x = 15 at.%.
Figure 8. TEM images of the microstructure and the corresponding SAED of alloys Ni49.5Ti50.5−xZrx: (a) x = 3 at.%, (b) x = 10 at.% and (c,d) x = 15 at.%.
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Figure 9. TEM images of the microstructure of the alloy Ni50Ti32Hf18 (a) bright field and (b,c) dark-field images and (d) correspondent SAED.
Figure 9. TEM images of the microstructure of the alloy Ni50Ti32Hf18 (a) bright field and (b,c) dark-field images and (d) correspondent SAED.
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Figure 10. (a) SEM and (bf) TEM microstructure images of Ti50Ni40Cu10 alloy obtained (ad) at room temperature and (e,f) at 120 K.
Figure 10. (a) SEM and (bf) TEM microstructure images of Ti50Ni40Cu10 alloy obtained (ad) at room temperature and (e,f) at 120 K.
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Figure 11. SEM images of (a) the packet twinned microstructure and (b) magnetic domain microstructure of the alloy Ni50Mn28.5Ga21.5.
Figure 11. SEM images of (a) the packet twinned microstructure and (b) magnetic domain microstructure of the alloy Ni50Mn28.5Ga21.5.
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Figure 12. TEM images of the microstructure of the alloy Ni54Mn21Ga25 and related SAED (a,b) long-period 10 M martensite at room temperature and (c,d) long-period 14 M martensite at 130 K.
Figure 12. TEM images of the microstructure of the alloy Ni54Mn21Ga25 and related SAED (a,b) long-period 10 M martensite at room temperature and (c,d) long-period 14 M martensite at 130 K.
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Figure 13. TEM images and corresponding SAEDs of rapid quenching from melt (RQM) alloys (a,b) Ti50Ni25Cu25 and (c,d) Ni45Ti32Hf18Cu5 obtained in an amorphous state and subjected to various crystallization annealing.
Figure 13. TEM images and corresponding SAEDs of rapid quenching from melt (RQM) alloys (a,b) Ti50Ni25Cu25 and (c,d) Ni45Ti32Hf18Cu5 obtained in an amorphous state and subjected to various crystallization annealing.
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Figure 14. TEM images of alloy (ac) Ti49.4Ni50.6 after high-pressure torsion (HPT), 5 revs, at 7 GPa, and subsequent annealing 523 K, 20 min and (d) Ti49.8Ni50.2 after step equal-channel angular pressing (ECAP) at 773 K + 723 K + 673 K, six passes.
Figure 14. TEM images of alloy (ac) Ti49.4Ni50.6 after high-pressure torsion (HPT), 5 revs, at 7 GPa, and subsequent annealing 523 K, 20 min and (d) Ti49.8Ni50.2 after step equal-channel angular pressing (ECAP) at 773 K + 723 K + 673 K, six passes.
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Figure 15. Fractography of (a,b) alloy Ti49,4Ni50.6 (a): After quenching, (b): HPT five turns at 6 GPa and subsequent annealing at 823 K and (c) quenched alloy Cu69Al28Ni3.
Figure 15. Fractography of (a,b) alloy Ti49,4Ni50.6 (a): After quenching, (b): HPT five turns at 6 GPa and subsequent annealing at 823 K and (c) quenched alloy Cu69Al28Ni3.
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Table 1. Chemical composition, temperatures of thermoelastic martensitic transformations (TMTs) B2 ↔ R ↔ B19′, average number of valence electrons per atom ev/a and their concentration cv.
Table 1. Chemical composition, temperatures of thermoelastic martensitic transformations (TMTs) B2 ↔ R ↔ B19′, average number of valence electrons per atom ev/a and their concentration cv.
Alloy, at.%RB19′ev/acv
Ms′, KMf′, KAs′, KAf′, KMs, KMf, KAs, KAf, K
Ti50Ni50 3433193533737.000.280
Ti50Ni49Al13193053013232642452893096.930.279
Ti50Ni48Al22952802763012001782272536.860.278
Ti50Ni47Al32622472422701301051611886.790.277
Ti50Ni46Al4190179171198 6.720.275
Ti50Ni45Al5139131112142 6.690.274
Ni50Ti49Al13213013053242662472883096.990.281
Ni50Ti48Al23002722773031991622222516.980.281
Ni50Ti48,5Al2,5282253256285109 1792116.980.282
Ni50Ti47Al3260228232265 6.970.282
Ni50Ti46Al4173134147184 6.960.282
Ti50Ni49Mn13162902853212802372892996.970.279
Ti50Ni48Mn22952632593012081642302596.940.278
Ti50Ni47Mn3277242237282135891762026.910.277
Ti50Ni45Mn5237191189243 6.850.276
Ni50Ti49Mn13042662723052832352983207.030.281
Ni50Ti48,5Mn1,52962592663012471912652907.050.281
Ni50Ti48Mn22882532592951861222182407.060.282
Ni50Ti47Mn325321622326278 1121407.090.283
Ni50Ti46Mn4173134146183 7.120.283
Ti50Ni49Fe13072942983132782502893056.980.279
Ti50Ni48Fe22872662762972181892352586.960.279
Ti50Ni46Fe42642382522801551111862106.920.278
Ti50Ni45Fe5230202217240 6.900.277
Ti50Ni49Co13243053133303132653223236.990.280
Ti50Ni48Co23012832913112712162892906.980.279
Ti50Ni46Co42872662742972371752612676.960.279
Ti50Ni45Co52632352442721651012072206.950.279
Ti50Ni40Co10218183190226 6.900.277
Ti50Ni35Co151308089140 6.850.276
Table 2. Chemical composition, temperatures of TMTs B2 ↔ B19 ↔ B19′, number of valence electrons per atom ev/a and their concentration cv.
Table 2. Chemical composition, temperatures of TMTs B2 ↔ B19 ↔ B19′, number of valence electrons per atom ev/a and their concentration cv.
Alloy, at.%B19B19′ev/acv
Ms′, KMf′, KAs′, KAf′, KMs, KMf, KAs, KAf, K
Ti50Ni45Cu5 3152222553377.050.281
Ti50Ni42,5Cu7,5 3052092293187.080.282
Ti50Ni40Cu103123023073162802012222917.100.283
Ti50Ni35Cu153323173223372001811862077.150.284
Ti50Ni33Cu17337326331342 7.170.284
Ti50Ni30Cu20337327332345 7.200.285
Ti50Ni25Cu25340326333348 7.250.287
Ti50Ni22Cu28335323328342 7.280.288
Ti50Ni18Cu32331313322339 7.320.289
Ti50Ni16Cu34321307313330 7.340.290
Ti50Ni48Pd2 3303103353557.000.276
Ti50Ni45Pd5 3203003303507.000.27
Ti50Ni40Pd103303203303402702302402807.000.261
Ti50Ni25Pd25450410430470 7.000.237
Ti50Ni5Pd45710680710740 7.000.211
Ti50Ni48Pt2 3052803053307.000.270
Ti50Ni42Pt83002752803052702402602907.000.245
Ti50Ni25Pt25720670720770 7.000.194
Ti50Ni47Au3 3152803203507.030.265
Ti50Ni45Au5 3603053203757.050.256
Ti50Ni38Au123653403553803002552503107.120.229
Ti50Ni25Au25425365400460 7.250.192
Ti50Au50880850870900 7.500.149
Ni50Ti47Zr3 3503253603857.000.274
Ni50Ti45Zr5 3503253703957.000.27
Ni50Ti40Zr10 3753504004207.000.261
Ni50Ti35Zr15 4554254755007.000.253
Ni50Ti32Zr18 5104805355557.000.248
Ni50Ti30Zr20 5555205706007.000.245
Ni50Ti38Hf12 4053804304507.000.226
Ni50Ti35Hf15 4504304805007.000.215
Ni50Ti32Hf18 5004755305557.000.206
Ni50Ti30Hf20 5204955645857.000.200
The chemical composition of the basic elements varied within 0.1 at.% and O and C were within 0.07–0.1 wt.%.
Table 3. Chemical composition, temperatures of TMTs L21 ↔ 2M (14M), the number of valence electrons per atom ev/a and their concentration cv of Ni–Mn–Ga-based alloys.
Table 3. Chemical composition, temperatures of TMTs L21 ↔ 2M (14M), the number of valence electrons per atom ev/a and their concentration cv of Ni–Mn–Ga-based alloys.
Alloy, at.%Ms, KMf, KAs, KAf, Kev/acv
Ni50Mn5097092097010208.500.321
Ni50Mn48Ga29078809209548.420.316
Ni50Mn46Ga48608308608978.340.312
Ni50Mn44Ga68077858308458.260.308
Ni50Mn42Ga8755 7908.180.303
Ni50Mn40Ga10693 7208.100.299
Ni50Mn36Ga145465255405577.940.290
Ni50Mn34Ga164574504604697.860.286
Ni50Mn32Ga183783753753907.780.282
Ni50Mn31Ga193683633633707.740.280
Ni50Mn30Ga203623563703747.700.278
Ni50Mn29Ga213283223353437.660.276
Ni50Mn28.5Ga21.53123063233307.640.275
Ni50Mn28Ga222942902963017.620.274
Ni50Mn27.5Ga22.52882842902947.600.273
Ni50Mn27Ga232782732782837.580.272
Ni50Mn26Ga242232162242297.540.270
Ni50Mn25Ga252001852162347.500.268
Ni53Mn22Ga252952802863037.590.270
Ni54Mn21Ga253133093103147.620.271
Ni54.75Mn20.25Ga253253203233307.640.272
Ni55.25Mn19.75Ga253483463503527.660.272
Ni55.75Mn19.25Ga253683633643707.670.272
Ni56.75Mn18.25Ga253683643653747.700.273
Ni57.5Mn17.5Ga255355265285407.730.274
Ni58.25Mn16.75Ga256025965986107.750.274
Ni59Mn16Ga256236106256357.770.275
Ni60Mn15Ga256296156306407.790.276
The chemical composition of the basic elements varied within 0.1 at.% and O and C were in the range of 0.1 wt.%.
Table 4. As and Af of TMT, the number of valence electrons per atom ev/a, and their concentration cv, average grain size <d> and mechanical properties of Cu–Al–Ni-based alloys.
Table 4. As and Af of TMT, the number of valence electrons per atom ev/a, and their concentration cv, average grain size <d> and mechanical properties of Cu–Al–Ni-based alloys.
Alloy, at.%As, KAf, Kev/acv<d>, μmσm, MPaσu, MPaδ, %Ψ, %
Cu78Al19Ni39039439.450.36460280780150.5
Cu77Al20Ni38939339.330.36380260520100.5
Cu75Al22Ni37938339.210.36213026049060.5
Cu73Al24Ni36737239.050.36035026045050.5
Cu71Al26Ni35135608.890.35875020039040.5
Cu69Al28Ni32833338.770.357100012025030.5
The chemical composition of the basic elements varied within 0.1 at.% and O and C were within 0.07–0.1 wt.%.
Table 5. Chemical composition, temperatures of TMTs B2 ↔ B19′, the number of valence electrons per atom ev/a and their concentration cv.
Table 5. Chemical composition, temperatures of TMTs B2 ↔ B19′, the number of valence electrons per atom ev/a and their concentration cv.
Alloy, at.%B19′ev/acv
Ms, KMf, KAs, KAf, K
Ti50Ni49.5Cr0.5 [38] *2922643023296.980.279
Ti49Ni50Cr1 [38] *2372152522687.020.281
Ti49Ni50V1 [38] *3182843273487.010.280
Ti48Ni50V2 [38] *3092853163407.020.281
Ti47Ni50V3 [38] *2932713003187.030.281
Ti45Ni50V5 [38] *2832692862977.050.281
Ti44Ni50V6 [38] *2792612822957.060.282
Ti50Ni45Cu5 [38] *3453173403687.050.281
Ti50Ni40Cu10 [38] *3062853003167.100.283
Ti50Ni45Cu5 [43] **3092953113307.050.281
Ti50Ni40Cu10 [43] **2921791922007.100.283
The content of C and O was less than 0.01 wt.% (*) and the content of O and C was nearly 0.1 wt.% (**).
Table 6. The magnitudes of the volume effect (V/V) and lattice parameters of the B2 austenite and of the B19, and B19′ martensites near the Ms temperature of the binary and ternary TiNi-based alloys.
Table 6. The magnitudes of the volume effect (V/V) and lattice parameters of the B2 austenite and of the B19, and B19′ martensites near the Ms temperature of the binary and ternary TiNi-based alloys.
Alloy, at.%ΔV/V, %B2Martensite B19′ (B19)
a, nma, nmb, nmc, nmβ, °
Ti50Ni50−0.130.30150.28900.41200.463096.8
Ti49.4Ni50.6−0.310.30130.28760.41320.462297.0
Ti50Ni49Co1−0.240.30140.28820.41150.464497.3
Ti50Ni46Co4−0.860.30130.28740.41080.463097.2
Ti50Ni45Co50.490.30120.28710.41040.469797.1
Ti50Ni45Cu5 [46]−0.50.30270.29030.41470.461396.2
Ti50Ni45Cu7.5 [46]0.360.30300.29110.42280.451790
Ti50Ni45Cu7.5 [46]−0.270.30300.29070.41700.459695.5
Ti50Ni40Cu10 [46]−0.360.30310.29010.42490.451590
Ti50Ni40Cu10 [46]−1.440.30310.28720.41920.457795.2
Ti50Ni37.5Cu12.5 [46]−0.140.30340.28960.42520.451490
Ti50Ni35Cu15 [46]−0.440.30430.28990.42600.451690
Ti50Ni30Cu20 [46]−0.530.30460.29000.42640.451290
Ti50Ni41Pd9 [47]−0.780.30470.28460.43040.458390
Ti50Ni40Pd11 [47]−0.930.30500.28300.43140.460490
Ti50Ni32Pd18 [47]−0.680.30560.28200.43430.462890
Ti50Ni30Pd20 [47]−0.580.30510.28200.43400.461390
Ti50Ni25Pd25 [47]−0.530.30630.28070.43610.466790
Ti50Au50 [48]−0.110.32200.2940.4630.49090
Table 7. The martensite lattice parameters of ternary TiNi- and Ni2MnGa-based alloys.
Table 7. The martensite lattice parameters of ternary TiNi- and Ni2MnGa-based alloys.
Alloy, at.%a, nmb, nmc, nmβ, °
Ti50Ni5Pd450.28030.4540.479490
Ti50Ni48Pt20.28920.41350.464396.6
Ti50Ni42Pt80.28210.42920.458590
Ti50Ni42Pt80.28730.42110.462795.6
Ti50Ni25Pt250.27650.44830.474490
Ti50Ni47Au30.29070.41360.465396.6
Ti50Ni45Au50.29070.41590.466696.5
Ti50Ni38Au120.28760.43190.4620390
Ti50Ni25Au250.28640.44930.4749290
Ni50Ti47Zr30.2920.41140.46797.8
Ni50Ti45Zr50.2930.4110.47298.5
Ni50Ti40Zr100.2980.4100.478100.3
Ni50Ti35Zr150.3030.4090.487101.8
Ni50Ti32Zr180.3050.4080.492102.5
Ni50Ti30Zr200.3070.4080.495103.7
Ni50Ti38Hf120.30060.41160.4803101.2
Ni50Ti35Hf150.30250.40960.4826102
Ni50Ti32Hf180.30510.40900.4850102.5
Ni50Ti30Hf200.30630.40830.4890103
Ni50Mn25Ga250.5550.5550.67090
Ni53Mn22Ga250.41420.5562.95390
Ni54Mn21Ga250.42320.5502.93793.3
Ni54.75Mn20.25Ga250.5490.5490.647890
Ni55.25Mn19.75Ga250.5480.5480.65390
Ni56.75Mn18.25Ga250.5460.5460.65390
Ni57.5Mn17.5Ga250.5420.5420.66090
Ni58.25Mn16.75Ga250.5430.5430.66390
Ni59Mn16Ga250.5430.5430.66390
Ni59.75Mn15.25Ga250.5430.5430.66390
Table 8. Mechanical properties of fine-grained (FG) alloys at room temperature.
Table 8. Mechanical properties of fine-grained (FG) alloys at room temperature.
Alloy, at.%σm, MPaσy, MPaσu, MPaσr, MPaδ, %er, %
Ti50Ni502008501200650406
Ti50Ni48Cu2150670850520205
Ti50Ni45Cu5120580750460184
Ti50Ni40Cu10100530750430253
Ti50Ni35Cu1570490720420203
Ti50Ni30Cu2060500720440203
Ti50Ni25Cu2560500740440203
Ti50Ni49Fe12007201000520245
Ti50Ni47Fe3 650950 18
Ti50Ni49Co11207401100620185
Ti50Ni48Co21607401100580184
Ti50Ni47Co32407101100470163
Ti50Ni45Co5 6501090 15
Ti50Ni43Co7 6301080 14
Ti50Ni10Pd403808501010470115
The content of O and C in the range of 0.07–0.1 wt.%.
Table 9. Mechanical properties of high-purity alloys Ti49.4Ni50.6, quenched and subjected to ECAP (1), HPT (2) and annealing.
Table 9. Mechanical properties of high-purity alloys Ti49.4Ni50.6, quenched and subjected to ECAP (1), HPT (2) and annealing.
Alloy, at.%<d>, mσm, MPaσy, MPaσu, MPaσr, MPaδ, %εr, %
Ti49.4Ni50.6601306301600500755
Ti49.4Ni50.6 (1)0.4025012001600950608
Ti49.4Ni50.6 (2)0.034601800210013401610
Ti49.4Ni50.6 (2)0.053701500167011302313
Ti49.4Ni50.6 (2)0.10200112012509203016
Ti49.4Ni50.6 (2)0.6025090012106507516
*—impurity content, C—0.0372 wt.%, O—0.0167 wt.%, N—0.003 wt.%, S—0.0001 wt.%.
Table 10. Mechanical properties of RQM alloys of the Ti–Ni–Cu system.
Table 10. Mechanical properties of RQM alloys of the Ti–Ni–Cu system.
Alloy, at.%<d>, mσm, MPaσy, MPaσu, MPaσr, MPaδ, %εr, %
Ti50Ni25Cu251.060680850620125
Ti50Ni25Cu250.870720900650124
Ti50Ni25Cu250.58010701200990114
Ti50.5Ni25Cu24.50.3801050117097093
Ti50.5Ni24.5Cu250.390115013801060103
Ti49.5Ni25Cu25.50.3809501120870103
Ti49Ni25Cu260.290115013801060103
Ti51Ni24Cu250.210011601300106093
Ti51Ni25Cu240.290120015501110103
The chemical composition of the basic elements varied within 0.1 at.% and O and C were within 0.07–0.1 wt.%.

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