Physical and Structural Characterization of Monocrystalline Cu-13.7% Al-4.2% Ni alloy submitted to Thermo-Cyclical Treatments under Applied Loads

Monocrystalline alloy with a nominal composition of Cu-13.7% Al-4.2% Ni (wt.%) that shows reversible martensitic transformations (RMTs) was studied. The alloy, manufactured by the “Memory Crystals Group” in Russia, was subjected to thermo-cyclical treatment (TCT) under tension within a range that included critical RMT temperatures. A special device was developed to perform TCTs (up to 500 cycles) and three different loads were applied: 0.11, 0.26, and 0.53 MPa. X-ray diffraction analysis, optical microscopy, differential calorimetry, and Vickers microhardness were involved in the alloy’s characterization. Under TCTs, the alloy displayed complex structural transformation, revealing the sequence of RMT, β1 ↔ R ↔ β′1 + γ′1; the involved phases were coherently precipitated but very sensitive to the experimental conditions. It was found that during TCTs (from 300 cycles) performed under optimum load (0.26 MPa), the processes of martensite reorientation, hardening, and stabilization of the structure were the most intensive thus leading to a reduction of RMT critical intervals and increased microhardness.

Monocrystalline Cu-Al-Ni (Cu 3 Al-based) alloys have received special attention, because they display high SME parameters along with good thermoelastic properties, formability, and plasticity without embrittlement at high temperatures [7][8][9][10][11][12][13]. Such SMAs are based on the Cu 3 Al intermetallic compound; they have different structures, depending on the composition, temperature, heat treatment, occurs by cooling below M f , an intermediate structure, the R-phase, related to the γ 1 ↔ β 1 is found. This is apparently a consequence of accumulated structural defects that retard the RMT and tend to stabilize the intermediate stage. The alloy submitted to TCs under a load of 0.5 kg revealed a tendency towards a decrease in the critical intervals for both direct and reverse martensite transformation with an increasing number of cycles. This tendency is accentuated for treatments up to 300 cycles in which critical temperature intervals of RMTs smaller than 10 • C were recorded. The hysteresis of the RMT also decreased. This can be attributed to a higher participation of the R-phase, particularly above 200 cycles, either alone or coherent with γ 1 (direct on cooling) and β 1 (reverse on heating). It is also suggested that the applied load contributes to the shape deformation observed in blocks of martensite variants. Despite complex structural transformations, the monocrystalline Cu-13.5wt.% Al-4.0wt.% Ni alloy showed good resistance to irreversible changes during applied tests.
In a more recent work [21], we paid attention to the behavior of a monocrystalline Cu-13.7Al-4.2Ni alloy subjected to thermal cycling treatment induced for 1, 100, 200, 300, 400, and 500 cycles without applied load, i.e., stress free.
In the current work, which revisits and expands our research, monocrystalline Cu-13.7% Al-4.2% Ni alloy was subjected to up to 500 cycles under three constant loads within a critical RMT temperatures range along with the physical and structural characterization of the alloy during applied TCTs.

Materials and Methods
The basic material investigated was a high purity monocrystalline Cu-13.7wt.% Al-4.2wt.% Ni alloy, obtained as a cylindrical bar (4 mm in diameter) from the Memory Crystals Group of the Technical University of Saint Petersburg, Russia [23]. Specimens approximately 4 mm in thickness were sectioned out perpendicular to the bar axis by means of a Minitom cutter (Struers A/S Pederstrupvej, Denmark). For microstructure characterization, the specimens were polished to an acceptable mirror finish using 0.1 µm alumina paste. No chemical etching was used, since the surface relief of transformation was enough to reveal the phases.
Specimens were submitted to 1, 100, 200, 300, 400, and 500 thermal cycles of heating up to 90 • C (above A f -in the initial state) followed by cooling to 0 • C (below M f -in the initial state). Each TCT was carried under a static load (TCTL). Three loads were applied: 0.14, 0.35, and 0.70 kg, corresponding to the stress of 0.11, 0.26, and 0.53 MPa, respectively. The chosen loads were below the flow limit determined as 70 MPa for this alloy [18,21]. The load was applied along with thermocycling using the special device developed for this purpose [18][19][20][21][22]. After TCTL, the structure of the alloy was analyzed by optical metallography and X-ray diffraction. Structural characterization was conducted at room temperature (RT) by finishing the treatment: (a) with a half cooling cycle to 0 • C, then up to RT, and (b) with a half heating cycle to 90 • C, then down to RT [12,[18][19][20][21][22].
The phase structure of the alloy was determined at RT by X-ray diffraction (XRD) in both Shimadzu 7000 (SHIMADZU Corporation, Kyoto, Japan) and Seifert 65 diffractometers (Seifert GmbH, Ahrensburg, Germany) using Cu-K α and Co-K α radiations, respectively, for 2θ angles from 25 • to 75 • at scanning steps of 0.03 • /3 s.
The surface morphological relief induced by the phase transformation and the involved deformation-related defects were observed in a Neophot-32 microscope (Carl Zeiss AG, Oberkochen, Germany) with polarized light. The microhardness was measured with a Shimadzu HMV-2 Micro Vickers Hardness Tester (Shimadzu Corporation, Kyoto, Japan). The RMT critical temperatures and associated thermal effects after the TCTL were determined using a Q10 differential scanning calorimeter (TA Instrument Co. Ltd., New Castle, DE, USA). The thermal tests were conducted in the temperature range from −50 • C to +150 • C with a heating/cooling rate of 10 • C/min in an inert atmosphere. Cooling was carried out using an LNCA mode (liquid nitrogen cooling). After TCTL, the samples were submitted to two thermal cycles during the DSC runs. In the first cycle of the DSC analysis, the samples were cooled to −50 • C, then heated to +150 • C, and cooled and heated again within the indicated temperature range. To determine the thermal hysteresis value in our studies, we made use of the The DSC analysis of monocrystalline Cu-13.7% Al-4.2% Ni alloy in the initial state ( Figure 2) showed that on the second heating, the reverse RMT (β′1 + γ′1) → β1 occurred in the range from +38.2 °C (As) to +68.8 °C (Af) through an endothermic process with a phase transformation enthalpy of 9 J/g. The process developed an "on-set" mode from +51.4 °C (Ao), presenting the peak temperature of +54.7 °C (Ap), where the heat flux was maximum. The process ended in an "off-set" mode at the temperature of +60.4 °C (Aof). On cooling of the alloy, the direct RMT β1 → (β′1 + γ′1) occurs between +51.4 °C (Ms) and 16.0 °C (Mf), having on-set temperature peaks 47.7 °C (Mo), +42 °C (Mp) and +32.6 °C (Mof), respectively, attributed to the exothermic process with a transformation enthalpy of 8.8 J/g.
It is known that SME Cu-Al-Ni alloys can undergo a single transformation (β1 → β′1 or β1 → γ′1) or a mixed transformation (β1 → γ′1 + β′1) depending on the composition of the alloy. In [24][25][26][27][28][29] it was The XRD pattern of the alloy in the as-received state in Figure 1a shows a significant amount of the β 1 martensitic phase, determined by the intensive β 1 and β 1 peaks, traces of the R-phase, determined by the peak (107)R, and the other peak found for the R-phase coinciding with the peak belonging to the β 1 phase and β 1 |(0213)R.
The DSC analysis of monocrystalline Cu-13.7% Al-4.2% Ni alloy in the initial state ( Figure 2) showed that on the second heating, the reverse RMT (β 1 + γ 1 ) → β 1 occurred in the range from +38.2 • C (A s ) to +68.8 • C (A f ) through an endothermic process with a phase transformation enthalpy of 9 J/g. The process developed an "on-set" mode from +51.4 • C (A o ), presenting the peak temperature of +54.7 • C (A p ), where the heat flux was maximum. The process ended in an "off-set" mode at the temperature of +60.4 • C (A of ). On cooling of the alloy, the direct RMT β 1 → (β 1 + γ 1 ) occurs between +51.4 • C (M s ) and 16.0 • C (M f ), having on-set temperature peaks 47.7 • C (M o ), +42 • C (M p ) and +32.6 • C (M of ), respectively, attributed to the exothermic process with a transformation enthalpy of 8.8 J/g. narrow, approximately 30.6 °C for As-Af and 35.4 °C for Ms-Mf. As a measure of thermal hysteresis, the difference among the peak critical temperatures (ΔT = Ap − Mp) was determined at approximately 12.7 °C in the first thermal cycle [20][21][22]. According to References [3,[24][25][26][27][28][29], RMT features can be deduced from the magnitude of the thermal hysteresis which is in good agreement with the results obtained, whereas the martensitic transformation β1 ↔ (β′1 + γ′1) presented hysteresis around 10 °C (Figure 3).   It is known that SME Cu-Al-Ni alloys can undergo a single transformation (β 1 → β 1 or β 1 → γ 1 ) or a mixed transformation (β 1 → γ 1 + β 1 ) depending on the composition of the alloy. In [24][25][26][27][28][29] it was confirmed that the mixed transformation (β 1 → γ 1 + β 1 ) takes place in the structure of Cu-Al-Ni alloys of approximately the same composition as the studied alloy. These martensitic transformations to γ 1 or β 1 structures have negligible energy difference so that sliding along the atomic layers can be implemented. Thus, the transformation from one compact structure to another is easily accomplished [3,30].
The DSC results show that during heating/cooling of the alloy, the critical RMT ranges were narrow, approximately 30.6 • C for A s -A f and 35.4 • C for M s -M f . As a measure of thermal hysteresis, the difference among the peak critical temperatures (∆T = A p − M p ) was determined at approximately 12.7 • C in the first thermal cycle [20][21][22]. According to References [3,[24][25][26][27][28][29], RMT features can be deduced from the magnitude of the thermal hysteresis which is in good agreement with the results obtained, whereas the martensitic transformation β 1 ↔ (β 1 + γ 1 ) presented hysteresis around 10 • C ( Figure 3). confirmed that the mixed transformation (β1 → γ′1 + β′1) takes place in the structure of Cu-Al-Ni alloys of approximately the same composition as the studied alloy. These martensitic transformations to γ′1 or β′1 structures have negligible energy difference so that sliding along the atomic layers can be implemented. Thus, the transformation from one compact structure to another is easily accomplished [3,30].
The DSC results show that during heating/cooling of the alloy, the critical RMT ranges were narrow, approximately 30.6 °C for As-Af and 35.4 °C for Ms-Mf. As a measure of thermal hysteresis, the difference among the peak critical temperatures (ΔT = Ap − Mp) was determined at approximately 12.7 °C in the first thermal cycle [20][21][22]. According to References [3,[24][25][26][27][28][29], RMT features can be deduced from the magnitude of the thermal hysteresis which is in good agreement with the results obtained, whereas the martensitic transformation β1 ↔ (β′1 + γ′1) presented hysteresis around 10 °C ( Figure 3).    Thermal cycling-treated samples under loads were submitted to two thermal cycles in the DSC analysis. Firstly, the samples were cooled from RT. The reverse RMT curves obtained in the first DSC cycle of monocrystalline Cu-13.7% Al-4.2% Ni alloy submitted to TCT under load are shown in Figure 4. After TCT under loadings of 0.11 MPa (Figure 4a) and 0.53 MPa (Figure 4c), the reverse RMT peak was split into two. In contrast, at TCT under a load of 0.26 MPa, the reverse RMT curves of the alloy showed a single transformation peak; thus, structural imperfections accumulated during TCTs promote full RMT (Figure 4b). Thermal cycling-treated samples under loads were submitted to two thermal cycles in the DSC analysis. Firstly, the samples were cooled from RT. The reverse RMT curves obtained in the first DSC cycle of monocrystalline Cu-13.7% Al-4.2% Ni alloy submitted to TCT under load are shown in Figure  4. After TCT under loadings of 0.11 MPa (Figure 4a) and 0.53 MPa (Figure 4c), the reverse RMT peak was split into two. In contrast, at TCT under a load of 0.26 MPa, the reverse RMT curves of the alloy showed a single transformation peak; thus, structural imperfections accumulated during TCTs promote full RMT (Figure 4b).
This result becomes more evident when considering reverse RMT enthalpies (E) of the alloy after TCT under the loads of 0.11 (Figure 4d), 0.26 (Figure 4e), and 0.53 MPa (Figure 4f), where E1 and E2 are the enthalpy of the first and the second transformation peak, respectively; E1+2 is the sum of the enthalpies.
It is interesting to monitor the evolution of the enthalpies E1 and E2. The value of E2 tends to increase with an increase in the number of cycles up to 100 TCT (at 0.11 MPa) or up to 300 TCT (at 0.53 MPa) and decrease with a further increase in the number of TCTs under load. The value of enthalpy was entire in the case of TCTs under 0.26 MPa (Figure 4e). The total enthalpy of the reverse transformation, in all tests, varied only slightly with the number of cycles.  An increase in the critical temperature range of complete transformation (β′1 + γ′1) → β1 correlate with the accumulation of lattice distortions during TMR/TCT, possible structural imperfections, its decrease with "training", and the predominant orientation of martensitic lamellas. Apparently, the load at 0.11 MPa was not large enough for "training" work, while the load at 0.53 MPa promoted the lattice distortions and the accumulation of its imperfections and complicated the implementation of TMR for the bulk material in the temperature range adopted in the TCT. In other words, under 0.11  (Figure 4f), where E 1 and E 2 are the enthalpy of the first and the second transformation peak, respectively; E 1+2 is the sum of the enthalpies.
It is interesting to monitor the evolution of the enthalpies E 1 and E 2 . The value of E 2 tends to increase with an increase in the number of cycles up to 100 TCT (at 0.11 MPa) or up to 300 TCT (at 0.53 MPa) and decrease with a further increase in the number of TCTs under load. The value of enthalpy was entire in the case of TCTs under 0.26 MPa (Figure 4e). The total enthalpy of the reverse transformation, in all tests, varied only slightly with the number of cycles.
An increase in the critical temperature range of complete transformation (β 1 + γ 1 ) → β 1 correlate with the accumulation of lattice distortions during TMR/TCT, possible structural imperfections, its decrease with "training", and the predominant orientation of martensitic lamellas. Apparently, the load at 0.11 MPa was not large enough for "training" work, while the load at 0.53 MPa promoted the lattice distortions and the accumulation of its imperfections and complicated the implementation of TMR for the bulk material in the temperature range adopted in the TCT. In other words, under 0.11 MPa and 0.53 MPa, the TCT was realized in conditions of incomplete RMT cycles.
In these tests, a load at 0.26 MPa seemed to be sufficient to give rise to the reversible martensitic transformations during heating and cooling thus facilitating the "training" of the alloy during cyclic tests and maintaining a narrow range of TMR and did not stimulate the accumulation of defects in the alloy structure.

Direct and Reverse RMT of the Monocrystalline Cu-13.7% Al-4.2% Ni Alloy Curves after TCT under Loads
The DSC analysis of the alloy in the initial state showed that on heating, the reverse RMT β 1 → β 1 occurred in the range of +38.2 • C (A s ) to +68.8 • C (A f ) through an endothermic process: phase transformation enthalpy~9 J/g. The process developed from an on-set temperature of +51.4 • C (A o ) and was responsible for the peak at +54.7 • C (A p ), where the heat flux was maximum. The process ended in an "off-set" mode at +60.4 • C (A of ). On cooling, the direct RMT, β 1 → β 1 , occurred in the temperature range of +51. The critical temperatures of intensive reverse RMT, Ao, Ap, and Aof, moved slightly to lower values, up to 300 cycles, and tended to increase afterward. The intensive ranges of reverse RMT, Ao-Ap and Ao-Aof, during the first thermal cycle were determined as 3.3 and 9 °C, then decreased to 2.7 and 8.6 °C after 300 cycles, and finally increased to 11.7 and 17.5 °C with 500 cycles TCT ( Figure 7).     The difference among the peak critical temperatures (ΔT = Ap − Mp), as a measure of the thermal hysteresis amplitude, was determined as approximately 12.7 °C in the first thermal cycle of the investigated alloy. The changes in peak temperatures, Ap and Mp, after TCT under load, alter the magnitude of the thermal hysteresis as observed in Figure 8.  • C, then decreased to 2.7 and 8.6 • C after 300 cycles, and finally increased to 11.7 and 17.5 • C with 500 cycles TCT (Figure 7).
The difference among the peak critical temperatures (∆T = A p − M p ), as a measure of the thermal hysteresis amplitude, was determined as approximately 12.7 • C in the first thermal cycle of the investigated alloy. The changes in peak temperatures, A p and M p , after TCT under load, alter the magnitude of the thermal hysteresis as observed in Figure 8. The difference among the peak critical temperatures (ΔT = Ap − Mp), as a measure of the thermal hysteresis amplitude, was determined as approximately 12.7 °C in the first thermal cycle of the investigated alloy. The changes in peak temperatures, Ap and Mp, after TCT under load, alter the magnitude of the thermal hysteresis as observed in Figure 8. The Cu-13.7%Al-4.2%Ni alloy diffractograms after TCT, terminated with the ½ cooling cycle (under load) and the ½ heating cycle (without the load), are shown in Figure 9a,b. As can be seen, the phase compositions of these alloys differ significantly.
With the accumulation of 200 thermal cycles, the diffractogram shows minor changes in structure, reducing only the intensity of the coherence planes peak β 1 |(0213)R and the R-phase peak (107)R. The martensitic β 1 phase was still observed, indicating its temperature/thermal stability. After 300 TCTs, the accumulated distortions promoted changes in the martensitic structure, and the induced martensitic transformation β 1 → γ 1 was observed. This TCT treatment promotes the greater stability of the martensitic phase γ 1 , determined by (011)γ 1 and (212)γ 1 peaks. The plane of coherence between the martensitic phases β 1 and R, β 1 |(0213)R, display moderate intensity.
After 200 and 300 TCTs, the induced martensitic transformation β 1 → γ 1 was incomplete (Figure 9b). The energy difference among these martensitic structures was negligible, and the sliding along of the atomic layers could be implemented. Thus, the transformation from one close-packed structure to the other was easily performed, mostly during thermal cycling [3,30].

Optical Microscopy after TCT under Loading
The structure of Cu-13.7%Al-4.2%Ni alloy after TCTs under the load of 0.11 MPa is shown in Figures 12 and 13. In the initial state, the alloy demonstrated a surface structure composed of martensitic blocks: a central block and four peripheral blocks. No significant changes were detected by optical microscopy at the end of treatment and ½ cooling cycle (Figures 12a,c,e and 13a,c) or ½ heating cycle (Figures 12b,d,f and 13b,d).
After TCTs under the load of 0.11 MPa, two martensitic blocks of the peripheral region were eliminated (Figure 12a-f). With the accumulation of 400 and 500 cycles (Figure 13a-d), the martensitic lamellae were not as distinct as before, but it was possible to examine microscopic characteristics of fine needles of this β′1 martensitic phase as estimated by XRD (Figure 9) [1,3].
The microstructure of Cu-13.7%Al-4.2%Ni alloy after TCTs under the load of 0.26 MPa is shown in Figures 14 and 15. The accumulation of structural imperfections during 100 and 200 cycles promoted the development of deep sliding lines, superimposed on the specific martensitic relief (Figure 14a,c). These lines originated from the multiple RMTs induced by TCTs [3,8,[18][19][20][21][22]. On heating, the superficial lines became microscopically visible due to the relief of tension (Figure 14b,d).

Optical Microscopy after TCT under Loading
The structure of Cu-13.7% Al-4.2% Ni alloy after TCTs under the load of 0.11 MPa is shown in Figures 12 and 13. In the initial state, the alloy demonstrated a surface structure composed of martensitic blocks: a central block and four peripheral blocks. No significant changes were detected by optical microscopy at the end of treatment and 1 /2 cooling cycle (Figure 12a changed; thus, differently oriented lamellae could be distinguished within the initial martensitic blocks (Figures 16e,f and 17a,d). After TCTs, the individual blocks of the martensitic lamellae with different orientations were observed throughout the sample's cross-sectional area. It is noted that the sliding lines caused by the RMT overlapped in certain regions, where the accumulation of structural defects probably occurs during TCT tests. Nevertheless, it can be said that the structural orientation was the result of the training under the applied load.  After TCTs under the load of 0.11 MPa, two martensitic blocks of the peripheral region were eliminated (Figure 12a-f). With the accumulation of 400 and 500 cycles (Figure 13a-d), the martensitic lamellae were not as distinct as before, but it was possible to examine microscopic characteristics of fine needles of this β 1 martensitic phase as estimated by XRD ( Figure 9) [1,3].
The microstructure of Cu-13.7% Al-4.2% Ni alloy after TCTs under the load of 0.26 MPa is shown in Figures 14 and 15. The accumulation of structural imperfections during 100 and 200 cycles promoted the development of deep sliding lines, superimposed on the specific martensitic relief (Figure 14a,c). These lines originated from the multiple RMTs induced by TCTs [3,8,[18][19][20][21][22]. On heating, the superficial lines became microscopically visible due to the relief of tension (Figure 14b,d). In the peripheral part a spear-shaped needle feature, martensite γ 1 (highlighted in Figure 14b) was presented prominently in the XRD pattern of Figure 10b. With the accumulation of 300 thermal cycles, the reorientation of the martensite lamellae could be noticed (Figure 14e,f). The microstructure of the alloy after TCTs of 400 and 500 cycles was practically stable, with the completely oriented martensitic lamellae in the central region (Figure 15a-d).   The morphological aspect of Cu-13.7% Al-4.2% Ni alloy after TCT under the load of 0.53 MPa is shown in Figures 16 and 17. With the accumulation of thermal cycles, the martensitic relief was less pronounced, probably due to the coherency between the martensitic β 1 and R-phases as noted in the diffractograms of Figure 11. With the accumulation of 300, 400, and 500 cycles, the martensitic relief changed; thus, differently oriented lamellae could be distinguished within the initial martensitic blocks (Figure 16e

Vickers Microhardness after TCT under Loads
The average values of Vickers microhardness as a function of the number of thermal cycles under the loads of 0.11, 0.26, and 0.53 MPa applied to the monocrystalline Cu-13.7% Al-4.2% Ni alloy are shown in Figure 18. Determined as 303 ± 7.8 kgf/mm 2 in the initial state, microhardness tended to decrease after TCTs under the loads of 0.11 and 0.53 MPa, going down to 241.8 ± 8.1 kgf/mm 2 and 268.7 ± 6 kgf/mm 2 , respectively, after 500 TCTs. However, after TCTs under the load of 0.26 MPa, microhardness displayed a stability plateau up to 300 thermal cycles. With the higher number of cycles, the microhardness increased from 295.3 ± 7.9 kgf/mm 2 , after 300 cycles, to 341.9 ± 14.5 kgf /mm 2 , after 500 thermal cycles, which could be typical for the more strained condition of any alloy with a stable structure due to the accumulation of imperfections that can hinder the development of the transformation. But, the apparent increase in microhardness can be associated with the recovery during unloading of the reversible strain accumulated by the induced martensitic transformation [1][2][3].

Vickers Microhardness after TCT under Loads
The average values of Vickers microhardness as a function of the number of thermal cycles under the loads of 0.11, 0.26, and 0.53 MPa applied to the monocrystalline Cu-13.7% Al-4.2% Ni alloy are shown in Figure 18. Determined as 303 ± 7.8 kgf/mm 2 in the initial state, microhardness tended to decrease after TCTs under the loads of 0.11 and 0.53 MPa, going down to 241.8 ± 8.1 kgf/mm 2 and 268.7 ± 6 kgf/mm 2 , respectively, after 500 TCTs. However, after TCTs under the load of 0.26 MPa, microhardness displayed a stability plateau up to 300 thermal cycles. With the higher number of cycles, the microhardness increased from 295.3 ± 7.9 kgf/mm 2 , after 300 cycles, to 341.9 ± 14.5 kgf /mm 2 , after 500 thermal cycles, which could be typical for the more strained condition of any alloy with a stable structure due to the accumulation of imperfections that can hinder the development of the transformation. But, the apparent increase in microhardness can be associated with the recovery during unloading of the reversible strain accumulated by the induced martensitic transformation [1][2][3].
/mm 2 , after 500 thermal cycles, which could be typical for the more strained condition of any alloy with a stable structure due to the accumulation of imperfections that can hinder the development of the transformation. But, the apparent increase in microhardness can be associated with the recovery during unloading of the reversible strain accumulated by the induced martensitic transformation [1][2][3].

Discussion
The investigated alloy demonstrated complex phase transformation involving coherent phases sensitive to experimental conditions. One of the key proofs of the coherence among the phases in the alloy structure was the absence of several DSC transformation peaks, contrary to that observed in similar alloys [3,[24][25][26][27][28][29]. A single RMT peak indicated the sequence of transformation β1 ↔ R ↔ β′1 + γ′1 in the structure of monocrystalline Cu-13.7% Al-4,2% Ni alloy.

Discussion
The investigated alloy demonstrated complex phase transformation involving coherent phases sensitive to experimental conditions. One of the key proofs of the coherence among the phases in the alloy structure was the absence of several DSC transformation peaks, contrary to that observed in similar alloys [3,[24][25][26][27][28][29]. A single RMT peak indicated the sequence of transformation β 1 ↔ R ↔ β 1 + γ 1 in the structure of monocrystalline Cu-13.7% Al-4,2% Ni alloy.
It should be noted that after TCTs finishing with the 1 2 heating cycle (AT → 90 • C → AT without loading), the XRD patterns of the alloys (Figures 9b, 10b and 11b) were in good agreement with the direct RMT critical temperatures (Figures 5a, 6a and 7a). Figure 5a shows that the temperature of the direct intensive RMT M of , determined after the first thermal cycle of the alloy, was above the TA, estimated by XRD analysis (20 ± 5 • C); it is assumed that this is within the martensite γ 1 field (below M of ) as observed in the diffractogram of Figure 9b. In general, when the critical temperature M of decreased approaching AT, a greater participation of the martensite β 1 and less participation of the others phases (γ 1 , β 1 |R and R) was observed in the diffractograms of Figures 9b, 10b and 11b.
It is noted that during all TCTs (Figures 5b, 6b and 7b), the XRD procedure for temperature (20 ± 5 • C) went beyond the reverse RMT temperatures even with the accumulation of TCTs, and the critical temperature A f was very far from TA. So, the participation of the high-temperature phase β 1 (mainly in the diffractograms after TCTs under the load of 0.26 MPa (Figure 10) was assumed to be a load response that promoted reverse RMT: γ 1 + β 1 → R → β 1 , only partially. These results are in agreement with our previous work [18][19][20][21][22], where the same TCTs of a Cu-Al-Ni alloy with comparable composition were performed; however, in TCTs without the applied load, it was not noticed.
During thermal cycling under load, two thermodynamic factors had a complex effect on the implementation of reversible martensitic transformations. Firstly, the load induced reversible MTs of the M → A def type, when heated from the martensitic phase temperature, or A → M def , when cooled from the austenitic phase temperature, thus orienting the structure in a certain way [1][2][3]32]. The fact of the possible reversible martensitic transformation of M-into A-phase under load, as predicted by Vasilevsky in 1971 [32], was experimentally proven in the studies on TiNi- [33][34][35][36][37] and Cu-Al-Ni [38]-based alloys with RMT.
Secondly, the load induced deformation that can be reversible (SME) or irreversible (inelastic deformation) in the alloys with RMT [2,3]. Reversible deformation depends on the load and it increases to a certain value (4-8%) initially but later on decreases. In this case, the temperature ranges for the realization of RMT and SMA expanded drastically as would be studied for a TiNi alloy [33][34][35][36][37].
According to the Clausius-Clapeyron equation, the applied external stress resulted in an increase in the critical temperatures of RMT [3,39]. On the other hand, the directed action of the load promoted some orientation of the structure and facilitated the transformation of martensitic lamellae according to certain variants [3,5].
As early as the 1980s in the works of Antipov et al. [40], the behavior of TiNi alloys with SME subjected to TCTs under different loads of flexural strain was analyzed. It was discovered that on heating during TCT, the return angle (determines the degree of shape restoration) decreased intensively in the first cycles and stabilized an increase in the number of thermal cycles, revealing the stabilization of the structure accompanied by a martensite reorientation and hardening. It was found that under optimized stresses, the return angle stabilization and stability of the structure develops with fewer cycles, while under low and/or excessive applied stresses, angle stabilization is delayed or even noticeable.
Comparison of the results of Antipov et al. [40] with the results of this work for the Cu-Al-Ni SMAs reveals a very similar tendency. With TCTs under the lowest applied load (0.11 MPa), structural stability with reorientation and hardening was not achieved even after TCT of 500 cycles, although critical temperatures and hysteresis showed fewer changes during cyclic tests as observed in Figures 5-8.
During TCTs under a higher applied load (0.26 MPa), RMT-induced processes of reorientation, hardening, and stabilization of the structure occurred more intensively, resulting in critical intervals reduction, microhardness increasing, and stabilization, from 300 thermal cycles. Such a conclusion was based on microhardness behavior, critical temperatures, hysteresis, and structural changes as observed in Figures 6, 8 and 18.
With the highest applied load (0.53 MPa), the reorientation and hardening processes could be expected to be faster; however, the excessive load acts both to facilitate the reorientation and to make the shape restoration difficult. As a result, the entire volume of alloy does not participate in the RMT in the limited temperature range adopted in the tests. In this process, the accumulation was more significant which slows down all the RMT reactions, increasing the critical intervals and revealing the intermediate structural states at the ambient temperature (TA) as observed in Figures 7 and 8.
Nevertheless, the changes in the main RMT parameters of monocrystalline Cu-13.7% Al-4.2% Ni alloy after TCT under all applied loads (after "training" under loads) were not very significant, especially after complete RMT cycles which indicates good resistance of the alloy investigated to irreversible structure changes.

1.
The monocrystalline Cu-13.7% Al-4.2% Ni alloy represents complex mixed transformation, revealing the RMT sequence β 1 ↔ R ↔ β 1 + γ 1 , where the present phases are coherent with each other and are very sensitive to structural changes.

2.
Alterations in the alloy structure by finishing the 1 /2 cycle heating treatment (TA → 90 • C → TA) show good consistency with the critical temperatures of direct RMT. The martensite γ 1 field lies below the end of the intensive direct RMT (below M of ). Above this temperature, a greater participation of martensite β 1 is observed, coherent to phase R. 3.
The participation of the high-temperature phase β 1 in the alloy structure, mainly after TCTs under the load of 0.26 MPa, appears as a response of the load applied during the treatment which promotes partially the reverse RMT γ 1 + β 1 → R → β 1 .

4.
During TCTs performed under optimized load (0.26 MPa), where the RMT range was not changed, the reorientation, hardening, and stabilization of the structure during RMT occur more intensively, resulting in a reduction of critical intervals, increasing and stabilizing of the microhardness, from 300 thermal cycles.

5.
The changes in the main RMT parameters of monocrystalline Cu-13.7% Al-4.2% Ni alloy during TCT under load are not very significant, especially after complete RMT cycles, which indicate good resistance of the alloy investigated to irreversible changes, making its practical use feasible.
Author Contributions: L.A.M.: conceptualization, project administration, creation of the special device for heat cycling, analyses of all results, original manuscript preparation; E.C.P.: investigation, the heat cycling tests performing, structure and DSC analyses, original figures preparation, discussion; S.A.P.: support and manufacture of the investigated alloy; C.Y.S.: DSC performing and DSC analysis, N.A.P.: discussion, translation into English and article edition, provision of referenced papers, including links and DOI. All authors have read and agreed to the published version of the manuscript.

Funding:
The researchers are grateful for the research support granted by FAPERJ, CAPES, and CNPq; N.A.P. thanks the support of state assignment No. 075-00947-20-00.