*3.4. EBSD and ODFs Results*

The recrystallization texture of the 97% cold-rolled samples after undergoing solution treatments at 1277 ◦C for 1 h was obtained using EBSD. Figure 7 shows a quasi-colored orientation map of the 97% cold-rolled samples after solution treatment, where the color indicates the crystal orientation presented in the stereographic triangle. In the figure, the RD, TD and ND represent the rolling direction, transverse direction and normal direction, respectively. From the results of orientation maps, the average grain size is around 402 μm. In the FeNiCoAlTaB alloy, the average grain size is around 400 μm [8]. From the inverse pole figure result, strong <001> orientation is observed in the rolling direction. The intensity of texture is 13.64. The recrystallization texture of the present alloy is similar to the <100> texture of the FeNiCoAlTaB and FeNiCoAlTiB alloy [8,10] but different from the <110> texture of the Fe–Ni–Co–Al–Nb–B alloy [9].

**Figure 7.** Quasi-colored orientation maps of microstructure and inverse pole figures of the cold-rolled alloy after being solution treated at 1277 ◦C for 1 h (RD: rolling direction, TD: transverse direction, ND: normal direction).

Figure 8 shows ODFs (phi 2= 0◦, 45◦, and 60◦ sections) for the alloy with different thermo-mechanical processing conditions. The texture of the as-received sample is cubic {001}100, Goss/brass (G/B) {110}115 and brass ({110}112. The sample after solution treatment shows cubic {001}100, rotated cube (Rt-C) {001}110 and rotated Cu (Rt-Cu) {112}011 texture. The 97% cold-rolled sample shows Goss (G) {110}001) and Goss/brass (G/B) {110}115 textures. The cold-rolled sample that underwent 1277C–1h heat treatment shows a Goss (G) {110}001 texture. Therefore, cold-rolling deformation and recrystallization at 1277 for 1h give rise to a preferred 100 recrystallized orientation.

**Figure 8.** ODFs results of FeNiCoAlTiNb polycrystalline under different thermo-mechanical processing conditions: (**a**) asreceived, (**b**) 1277C–24 h, (**c**) 97%CR and (**d**) 97%CR + 1277 ◦C for 1 h.

#### *3.5. Three-Point Bending Test Results*

The results from the thermal cycling under the three-point bending test at different stress levels are displayed in Figure 9. Figure 9a shows the thermal cycling results of sample aged at 600 ◦C for 24 h. When the stress level is 100 MPa, the martensitic transformation occurs at around −150 ◦C. The complete loop of phase transformation does not observe because the transformation temperature is too low, and the device can only cool down to −150 ◦C, and the stress level is not high enough to induce martensitic transformation to higher temperature. The sample fails during the 200 MPa thermal cycle test. The early failure of thermal cycling specimens was probably caused by the variant–variant interactions during transformation process. During the cooling process, the interaction between the growing martensitic phase increases variant interaction and leads to fracture [20].

Figure 9b presents the thermal cycling results of the sample aged at 600 ◦C for 48 h. The phase transformation was observed when the stress level was 50 MPa. The sample failed when the stress level increased to 250 MPa. Transformation temperatures, recoverable strain and temperature hysteresis values can be obtained through this test. The maximum recoverable strain was around 1.6% at a stress level of 200 MPa. The transformation temperatures (Af and Ms) are extracted in Figure 9b and plotted as a function of applied stress, as shown in Figure 9c. From the results, when the applied stress levels increased, the Af and Ms slightly increased their temperatures. The Ms was −110 ◦C at 50 MPa and increased to −100 ◦C at 200 MPa. The Af was −70 ◦C at 50 MPa and increased to −56 ◦C at 200 MPa. In addition, the lines for each transformation temperature could be extrapolated down to zero stress to determine the stress-free transformation temperatures. The Ms was

−112 ◦C and Af was −74 ◦C, which were values close to the transformation temperatures determined by the SQUID results.

**Figure 9.** Shape memory effect of FeNiCoAlTiNb 97% cold-rolled sample after being aged at 600 ◦C for x hours using three-point bending test. (**a**) 600 ◦C–24 h aging sample, (**b**) 600 ◦C–48 h aging sample, (**c**) stress versus temperature phase diagram for a 600 ◦C–48 h aging sample. The values for each transformation temperature are obtained from Figure 7b results. (**d**) 600 ◦C–72 h aging sample.

Figure 9d displays the thermal cycling results of the sample aged at 600 ◦C for 72 h. The maximum recoverable strain was 1.1%. The sample failed during the 150 MPa thermal cycle test. The sample was brittle due to the formation of β phases at the grain boundaries.

#### **4. Discussion**

From the three-point bending test results, it is clear that the recoverable strain of the 600 ◦C–48 h aged sample is smaller than the theoretical calculated values of FeNi-CoAlTaB [8]. One possible reason for this is the smaller fraction of low angle grain boundary (LAB). Figure 10 shows the grain boundary misorientation of the 97% cold-rolled alloy after solution treatment at 1277 ◦C for 1 h. Based on the calculation, the fraction of LABs is about 11%. The FeNiCoAXB (X: Ta, Nb, Ti) polycrystalline alloy systems are required to undergo aging heat treatment to increase the strength of the austenite matrix and obtain the nano size L12 precipitates [8–10,25]. Meanwhile, β phases (grain boundary precipitates) are generated along the grain boundary, especially at the triple junction. β phases prefer to form at high angle grain boundaries (HABs) due to higher energy [9,10]. From Figure 3c results, it is seen that the β phases form at the grain boundaries. As the fraction of LABs is 11% in the 97% cold-rolled sample after solution treatment at 1277 ◦C for 1 h, β phases can be seen in the grain boundaries for the 600 ◦C–48 h sample. β phases deteriorate ductility and shape memory properties. As a consequence, the recoverable strain of the

shape memory effect is smaller than the theoretical calculated values. Although a small number of β phases is formed in the 600 ◦C–24 sample, the transformation curve of the shape memory effect is not observed even under the 2000 MPa stress level due to its lowest transformation temperatures. For the 600 ◦C–72 sample, more β phases were generated in this aging condition, as shown in Figure 3d. As a consequence, the sample failed during the 150 MPa test.

**Figure 10.** Grain boundary misorientation of the cold-rolled alloy after solution treatment.

Another possible reason for the significantly lower shape memory recoverable strain observed in this study is related to the texture intensity of the <100> orientation in the rolling direction. In this study, the texture intensity of the <100> component in the RD direction is 13.64, as shown in Figure 7. The texture intensity of same component and orientation in FeNiCoAlTaB is 25.4 [8]. The value of texture intensity in this sample is lower than that in the FeNiCoAlTaB specimen. Texture development can improve the recoverable strain of shape memory and superelasticity [8–10]. In iron-based SMAs, reducing the grain constraint is crucial due to martensite variant selection. Based on the Bain distortion theory, only three variants can assist the deformation, which is different from the NiTi SMAs, in which twenty-four variants are able to assist deformation [1].

In iron-based SMAs, the theoretical transformation strains are 8.7%, 4.1% and 2.1% in <100>, <110> and <111> orientations, respectively [8]. Based on the results in Figure 7, Quasi-colored orientation maps in the rolling direction show that some grains are oriented in the <110> and <111> direction. The large difference in the crystal orientation between surrounding grains increases the grain constraint. When the applied stress level is increased, the incompatibility between adjacent grains becomes intense and results in the deterioration of ductility. Figure 11a presents SEM images of the fracture surface for the 600 ◦C–48 h sample after three-point bending test. The fracture is along the grain boundary. Figure 11b shows an SEM image of a cross section. The fracture surfaces of the sample are smooth. The results suggest that the sample has a brittle intergranular fracture, and the fracture is along a grain boundary.

The final possible reason for the smaller strain of this polycrystalline material is the addition of boron. The volume fraction of LABs is 11% in the present alloy and 11% in FeNiCoAlNbB [9]. Since both alloys have a similar fraction of LABs, the aging condition for the FeNiCoAlNbB alloy is 600 ◦C for 96 h. A small amount of grain boundary precipitates is observed in this aging condition. Since B element tends to lower the grain boundary energy, this retards β phases' growth in the FeNiCoAlNbB alloy. The recrystallization texture of the FeNiCoAlNbB alloy near <110> in the rolling direction is different from the FeNiCoAlTaB [8,9] and present alloy.

(**b**)

**Figure 11.** SEM images of 600 ◦C–48 h sample after three-point bending test. (**a**) Fracture surface of sample and (**b**) cross section of fracture surface.

#### **5. Conclusions**

In summary, the microstructure, thermo-magnetization, and shape memory properties of new cold-rolled FeNiCoAlTiNb SMA samples were investigated for the first time. The main conclusions can be listed as follows:


increased from 24 to 72 h, the martensitic transformation temperature shifted to a higher temperature, which indicates the transformation temperatures increase with an increase in the aging times.


**Author Contributions:** Conceptualization, L.-W.T.; methodology, L.-W.T. and C.-H.C.; validation, L.-W.T. and C.-H.C.; investigation, L.-W.T., C.-H.C., W.-C.C., Y.C. and N.-H.L.; resources, L.-W.T. and C.-H.C.; data curation, W.-C.C., Y.C. and N.-H.L.; writing—original draft preparation, L.-W.T.; writing—review and editing, L.-W.T.; visualization, L.-W.T. and C.-H.C.; supervision, L.-W.T.; project administration, L.-W.T. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Ministry of Science and Technology (MOST), grant numbers MOST 109-2221-E-018-010-MY2. This work was financially supported by the Young Scholar Fellowship Program of the Ministry of Science and Technology (MOST) in Taiwan, grant number MOST 110-2636-E-002-005.

**Acknowledgments:** The FeNiCoAlTiNb alloys were fabricated by the National Chung-Shan Institute of Science and Technology (NCSIST), which is gratefully acknowledged. The authors would like to thank Yung-Sheng Chen at the Instrumentation Center, National Tsing Hua University for the SQUID measurements. Thanks Ko-Kai Tseng at the High Entropy Materials Center, National Tsing Hua University for the cold rolling experiment.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


## *Communication*
