Preparation and Characterization of Ni-Mn-Ga-Cu Shape Memory Alloy with Micron-Scale Pores

Abstract

Porous Ni-Mn-Ga shape memory alloys (SMAs) were prepared by powder metallurgy using NaCl as a pore-forming agent with an average pore size of 20–30 μm. The microstructure, phase transformation, superelasticity, and elastocaloric properties of the porous alloys were investigated. The prepared porous alloy had a uniform pore distribution and interconnected microchannels were formed. Cu doping can effectively improve the toughness of a porous alloy, thus improving the superelasticity. It was found that porous Ni-Mn-Ga-Cu SMAs have a flat stress plateau, which exhibits a maximum elongation of 5% with partially recoverable strain and a critical stress for martensite transformation as low as about 160 MPa. In addition, an adiabatic temperature change of 0.6 K was obtained for the prepared porous alloy at a strain of 1.2% at about 150 MPa. This work confirms that the introduction of porous structures into polycrystalline Ni-Mn-Ga SMAs is an effective way to reduce costs and improve performance, and provides opportunities for engineering applications.

1. Introduction

Shape memory alloys (SMAs), represented by Ni-Mn-based alloys, have many unique properties such as a shape memory effect, superelastic effect, magnetocaloric effect, and elastocaloric effect [1], which make them useful for applications in refrigeration, shock absorption, sensors, and actuators [2]. The reversible martensite transformation process, stimulated by an external stress or magnetic field, is the fundamental result of the properties mentioned above. Due to the grain boundary constraints in polycrystalline alloys, phase transformation and the movement of martensitic variants are suppressed [3]. It has been shown that when the characteristic length (which is related to the mechanics of a phenomenon) interacts with a size parameter (e.g., grain size, film thickness, fiber diameter), the grain boundary constraints in that direction can be reduced substantially, resulting in an increase in the alloy’s physical properties, even coming close to those of a single-crystal counterpart [4]. A one-dimensional Ni-Mn-Ga microwire prepared using the glass coating method shows a bamboo-like structure, which can achieve up to 14% fully recoverable strain at room temperature [5,6]. Ni-Mn-Sn and Ni-Mn-Ga films acquired through DC magnetron sputtering have the characteristics of a “panpipe structure” and a maximum magnetic entropy change as high as 4.184 J/g·K [7]. Although low-dimensional SMAs have achieved high performance, they still face the challenge of complex preparation processes, a high brittleness, and size limitations.

In order to reduce the grain boundary constraints in polycrystalline alloys, a porosity strategy was proposed, where the pore walls in a porous alloy can be regarded as films, while the pillars behave like wires. Pores significantly decrease the grain boundary constraint and thus reduce the stress required to drive the martensite transformation [4]. In addition, pores provide a large amount of free surface, which allows fluid to flow through to improve the heat exchange efficiency, thereby improving the application of porous SMAs in magnetocaloric and elastocaloric refrigeration devices [8]. It has been reported that a porous Ni-Mn-Ga alloy with a porosity of 57% and pore size of 355–500 μm was prepared using the infiltration method [9]. By fabricating a dual-pore structure with both large and small pores, the researchers further reduced the number of grain boundaries [10]. The bottom-up powder metallurgy method offers greater flexibility when designing the size, shape, and distribution of the pores in porous alloys compared to other preparation processes [11]. A porous Ni-Mn-Ga alloy prepared by spark plasma sintering (SPS) using NaCl as the pore-forming agent has a high porosity of 92% [12]. There are also reports of the use of Mg powder as a pore space holder to prepare Ni-Mn-Ga porous alloys using one-step sintering [13]. Element doping is widely used in Ni-Mn-based alloys, such as Co [14], Fe [15], and Zn [16], to solve the intrinsic brittleness of Ni-Mn-based SMAs. And the addition of a fourth element causes the formation of precipitates at the grain boundaries in SMAs, which enhance their strength [17]. By doping with Cu, the grain boundaries were strengthened and a dual-phase structure containing martensite and the precipitated phase was observed. After heat treatment, the Ni₃₀Cu₂Mn_42.4Ga_7.6 alloy exhibits a dual-phase structure with a small amount of single-crystal γ-phase distributed along the boundaries of martensitic grains, effectively hardening the boundaries of the martensitic grains [18]. Additionally, researchers observed a dual-phase structure containing a martensitic phase in Ni₅₀Mn₂₅Ga₁₆Cu₉ alloys. The addition of Cu not only strengthens the grain boundaries, but also changes the martensitic structure, effectively improving the ductility of Ni-Mn-Ga alloys [19]. Moreover, the addition of Cu can lead to grain refinement, which can be explained by different crystallization processes [20]. And the addition of the Cu element into the Ni-Mn-In SMA can induce a greater heat absorption/release effect [21].

Ideal mechanical properties can broaden the application of Ni-Mn-based alloys. At present, most of the porous SMAs prepared by researchers have pore sizes over 100 μm [22]. However, smaller pores and a higher porosity will make the pore distribution more uniform and make it easier to obtain bamboo-like pillars. In our previous investigation, porous Ni-Mn-Ga alloys with a uniform distribution of pores with sizes of 20–30 μm were successfully fabricated by powder metallurgy. However, the poor strength resulted in a low performance [23]. In the current study, the microstructure, phase transition, mechanical properties, and elastocaloric effect of a Cu-doped Ni-Mn-Ga porous alloy are analyzed and discussed, which will provide theoretical support for further applications.

2. Experimental Procedures

Bulk ingots with a nominal composition of Ni₅₀Mn_25−xGa₂₅Cu_x (where x = 0, 2, 3 and referred to as Cu0, Cu2, Cu3) were produced by vacuum induction melting of a high-purity (99.99%) Ni-Mn alloy (to prevent the oxidation of the Mn element) and a Cu and Ga alloy in an argon atmosphere. The dimension of the ingots was 100 × 60 × 20 mm, and the weight was about 975 g. After being homogenized at 1173 K for 48 h, the ingots were crushed into a powder using a vibration grinder and sieved to a size of 20–30 μm. In order to prevent the oxidation of alloy powder during the grinding process, the grinding time should be controlled to be less than 30 s. NaCl powder was also sieved to 20–30 μm. The alloy powder and NaCl powder were mixed using a ball mill for 5 h. The volume ratio of the alloy powder to NaCl was 1:1 to achieve a 50% porosity. The mixed compacts were obtained through cold pressing at about 800 MPa. Finally, a vacuum tube furnace was used for high-temperature sintering at a vacuum value of up to 8 × 10⁻⁴ Pa. The sintering process included heating to 1083 K for 1 h to remove water vapor and NaCl (melting point: 1073 K), then heating to 1373 K for sintering for 2 h, and then pouring water onto the glass tube of the tube furnace for cooling. The phase transition temperature of the alloy was measured by differential scanning calorimetry (DSC, NETZSCH DSC 214 Polyma, Bavaria, Germany). The structure and phase analyses of the alloy were carried out by X-ray diffraction (XRD, Riguku SmartLab, Tokyo, Japan) with Cu Kα radiation (λ = 1.54 Å) at 292 K. The microstructure of the specimens was analyzed using scanning electron microscopy (SEM, Zeiss Supra 55, Zeiss, Oberkochen, Germany) and transmission electron microscopy (TEM, JEM-2200FS, Tokyo, Japan). The energy dispersive spectrometry (EDS) was used for composition analysis. The pore structure of the porous alloys was analyzed using X-ray microtomography (CT, YXLON FF35CT, Hamburg, Germany). The MTS 628 universal testing machine was used to test the mechanical properties of the specimens with dimensions of Φ3 × 6 mm, and the axial displacement of the specimens was monitored through an extensometer (YUU-25/5, NIM, Beijing, China). Due to the small size of the sample, the extensometer was fixed on the indenter of the universal testing machine. A K-type thermocouple was used to monitor the temperature change.

3. Results and Discussion

Figure 1a,b show the microstructure of the Ni-Mn-Ga and NaCl powder separately. Upon inspection, it can be seen that both powders have a size of about 20–30 μm. After grinding, the alloy powder and NaCl exhibit an irregular shape. Due to the sintering temperature (1373 K) being lower than the melting point of Ni-Mn-Ga (1407 K [24]), the alloy undergoes solid-phase sintering, which is conducive to the formation of sintering necks. As shown in Figure 1c, complete sintering necks were formed (red arrows), with uniform pore distribution. In addition, most of the pores are interconnected (blue arrows), creating more free surfaces. Three-dimensional tomography was performed using industrial CT, as shown in Figure 1d–f. The relative density of the porous alloy measured by industrial CT is 44%, slightly lower than the design (50%), which is due to the melting of the alloy surface during sintering, resulting in volume shrinkage. As shown in Figure 1d, due to the interconnected pores, the pores are confirmed as a single unit. However, there are still a small number of isolated pores (Figure 1e) with sizes smaller than 10 μm. Figure 1f illustrates the radial and axial porosity of the porous alloy, which is defined as the proportion of the total pore volume to the volume of the specimen. The blue area depicts the proportion of the alloy in that particular direction, whereas the green area signifies the pores. It can be seen that the pore distribution is uniform, which is beneficial for the uniform stress distribution of porous alloy pillars during compression.

The composition of the porous alloys measured by EDS is listed in

Table 1. The composition, density, and porosity of Cu0-3 bulk and porous alloys.

Specimens	Composition (%)								Density (g/cm3)		Porosity (%)
	Bulk				Porous				$\mathsf{\rho }$0	$\mathsf{\rho }$
	Ni	Mn	Ga	Cu	Ni	Mn	Ga	Cu
Cu0	49.1	24.5	26.3	0	50.2	23.9	25.8	0	8.1	4.4	45.2
Cu2	49.1	24.9	23.4	2.5	50.8	24.2	22.6	2.3	8.1	4.3	46.4
Cu3	50.0	23.4	23.1	3.4	51.4	22.1	23.3	3.2	8.2	4.5	45.2

. It can be seen that the matrix of the porous alloy is approximately the same as that of the bulk alloy. As a result of the specimen undergoing multiple heat treatment processes, there is a loss of the Ga element. Table 1 also includes the densities of both the bulk and porous alloys, along with the porosity of the porous alloys determined through the relative density method. The porosity of the porous alloys was obtained through the formula $\mathrm{P}=(1-\mathsf{\rho }/{\mathsf{\rho }}_0)\times 100\mathrm{\%}$, where P is the porosity, $\mathsf{\rho }$ is the density of the porous alloy, and ${\mathsf{\rho }}_0$ is the density of the bulk Ni-Mn-Ga alloy. The ${\mathsf{\rho }}_0$ was measured using the Archimedes drainage method, and $\mathsf{\rho }$ was obtained by cutting the porous alloy into standard cylinders. The porosity obtained through the relative density method is approximately 45%, which is equivalent to the density measured by industrial CT.

Figure 2 shows the XRD patterns of the Cu0, Cu2, and Cu3 bulks, as-milled powder, and porous alloys under room temperature (300 K). As shown in Figure 2a, the main phase of the bulk alloys is a typical L2₁ cubic austenitic structure. Upon the addition of Cu, the Cu2 specimen exhibits weak peaks of NM martensite at 78°, attributed to the enhancement of the e/a ratio of the material due to the Cu incorporation, which subsequently elevates the martensite transition temperature [25]. However, due to the internal stress generated during grinding, the diffraction peak of the powder specimen becomes wider (Figure 2b), which may submerge the peaks associated with martensitic phases. The diffraction peak width of those porous alloys becomes narrower, as shown in Figure 2c, which means the release of internal stress during sintering. The martensite diffraction peak can be clearly observed in the Cu2 and Cu3 specimens. The martensite peaks in the Cu2 and Cu3 specimens represent the coexistence of NM martensite with the I4/mmm space group [26] and 7M modulation martensite structure [27]. The internal stress generated during powder grinding and green billet pressing is not completely released, and the internal stress may also arise during cooling after sintering, ultimately leading to stress-induced martensitic transformation in certain grains.

DSC curves of the Cu0–3 alloys with bulk, powder, and porous structure are shown in Figure 2d,e. In Figure 2d, with the addition of Cu, the martensite transformation temperature of the bulk specimens gradually increases. Since the finishing temperature of martensitic transformation (M_f) for the Cu2 and Cu3 bulk alloys is close to the testing temperature, it may result in the presence of residual martensite phases, which aligns with the observed XRD pattern. In Figure 2e, the grinding process induces an uneven and non-directional internal stress state within the powder sample. This internal stress is caused by point defects (vacancies and self-interstitials), resulting from atomic disorder and severe plastic deformation, which can lead to a decrease in the stability of the martensitic phase, causing the disappearance of phase transition peaks [28]. In Figure 2f, the transformation peaks of the porous alloys exhibit a broader and less intense profile compared to those of the bulk alloys, which may be attributed to the internal stress generated during green billet pressing and sintering processes. Notably, the phase transformation temperature of the porous alloys is higher than that of the bulk alloys, owing to the Mn loss during sintering, which subsequently alters the e/a ratio.

The stress–strain curves of the Cu0, Cu2, and Cu3 porous alloys are shown in Figure 3a–d. The specimens were subjected to cyclic compression with a loading and unloading rate of 0.003 s⁻¹. For enhanced visibility, the curves were shifted along the X-axis. From the 2.5% stress–strain curve of the Cu0 specimen (Figure 3a), the specimen first undergoes elastic deformation of austenite. There is a gentle slope change of the curve, indicating the occurrence of stress-induced martensitic transformation (SIMT). It can be seen from Figure 3b that with the addition of Cu, the maximum strain increased from 2.5% of Cu0 to 4% of Cu2, and the maximum stress increased from 80 MPa to 160 MPa. Meanwhile, the maximum strain of the porous Ni-Mn-Ga alloy obtained with Mg as a pore-forming agent is about 2% [29]. Porous Ni-Fe-Ga alloys with pore sizes of 355–500 µm prepared by replication casting have a total recoverable strain of 3.98% under the stress of 60 MPa [30]. Due to the structural defects caused by pores, coupled with the intrinsic brittleness, most of the porous shape memory alloys have limited ductility.

It is noteworthy that when the strain exceeds 2.5%, a significant inflection point appears in the stress–strain curve, followed by a stress plateau. Due to the introduction of pores, the stress transmission path is limited, leading to stress concentration. And this part of the grains first started SIMT. As the stress increases, other grains progressively attain the critical stress for stress-induced martensitic transformation (σ_SIMT) and initiate the transformation process. During unloading, the reverse transformation of martensite and elastic recovery occur simultaneously. The presence of irreversible strain may be due to the generation of residual martensite. In addition, the fracture of the pillar caused by stress concentration also constitutes irreversible strain. For comparison, the bulk alloy did not undergo SIMT until 2% pre-strain (Figure 3d). However, due to the high brittleness, the specimen collapsed during 2.5% compression. It is noteworthy that the σ_SIMT of the bulk alloys reaches 280 MPa, which is significantly higher than that of the Cu2 porous alloys (160 MPa). The pores reduce the number of grain boundaries and thus reduce the phase transformation resistance. Lower stress is advantageous for minimizing crack initiation, and the adaptability of stress conduction in porous alloys further enhances their ductility [31].

To provide further clarification of the deformation process, the stress–strain curves of the recovery stages are decomposed and analyzed. Figure 3e shows the division of different strains, where ε_t, ε_irr, and ε_e are the total strain, irrecoverable strain, and elastic strain. And the ε_se (superelastic strain) is defined as ${\mathsf{\epsilon }}_{\mathrm{s}\mathrm{e}}={\mathsf{\epsilon }}_{\mathrm{t}}-{\mathsf{\epsilon }}_{\mathrm{e}}-{\mathsf{\epsilon }}_{\mathrm{i}\mathrm{r}\mathrm{r}}$. During unloading, elastic recovery first occurs, accompanied by the reverse transformation of martensite. Subsequently, irrecoverable strain is generated. With the increase in pre-strain, the proportion of ε_e decreases and ε_irr increases, which means the specimen mainly undergoes elastic strain at low pre-strain (Figure 3f). As the pre-strain increases, the increase in ε_irr becomes predominant. It is worth noting that the proportion of ε_se in the Cu2 and Cu3 specimens is basically unchanged at about 1.2%, which is significantly lower compared to wire and single-orientation polycrystalline alloys [32,33], but slightly higher than that of non-oriented bulk Ni-Mn-Ga alloys [23]. The low ε_se of the porous alloys is attributed to the randomness of the grain orientation in the porous alloys, subjecting the grains to a complex stress environment. Stress can only be transmitted along a fixed path, which makes porous alloys prone to stress concentration. As the stress increases, some grains undergo elastic deformation, while others experience martensitic transformation or plastic deformation. And the grains in porous structures undergo the above deformation stages asynchronously. Additionally, the SIMT of each grain is similar, resulting in a constant proportion of superelastic deformation. Although there is the phenomenon of stress concentration in bulk alloys [34], the presence of pores amplifies it.

Furthermore, Figure 4 shows the morphology of the Cu2 specimen. As illustrated in Figure 4a, the porous alloy fails with the shear bands at 45°, indicating a uniform internal structure [35]. The fracture surface of porous alloys (Figure 4a) can be divided into intergranular fracture (blue box) and transgranular fracture (red box). Among them, the number of intergranular fractures accounts for the majority. The transgranular fracture exhibits a planar appearance, resulting from the fracture of the sintering neck (Figure 4b). Precipitates of about 2 μm can be observed at the fracture (red arrow) and free surface (blue arrow), which was analyzed by EDS to be Mn-O, while the component of the matrix is the same as the design (orange arrow). In Figure 4c, the transgranular fracture shows a predominant cleavage river pattern (red arrow), which is a typical brittle fracture. And no obvious Mn-O precipitation was found at the fracture surface, which means the Mn tends to precipitate on the free surface as well as along grain boundaries. However, the Mn segregation was not found in the alloy powder, indicating that Mn precipitates emerged during sintering. In Figure 4d, the martensitic lath spans the sintering neck. And the EDS line scanning at the sintering neck shows no obvious element segregation (Figure 4e). The aforementioned phenomenon demonstrates that the grains within the sintering neck have undergone effective sintering. It is apparent that no significant component segregation was detected at the fracture surface of the Cu0 specimen (Figure 4f). And the strength of the Cu0 specimen is significantly lower than that of the Cu2 specimen (Figure 3a). The above phenomena can lead to the conclusion that the Mn-rich particles at the intergranular fracture enhance the bonding between grains. During compression, the stress concentration causes the crack initiation at the grain boundaries and sintering neck, which means that intergranular fracture occurs first during fracture. The Mn-rich particles impede the movement of dislocations, thereby enhancing the bonding strength between grains. After the sintering neck fractures, due to the self-regulation of stress in the porous structure, the stress is concentrated on the other grains, causing transgranular fracture.

Notably, Mn precipitates with a diameter of about 150 nm were also found in polycrystalline Ni₂MnGa cylinders obtained by suction casting solidification and melt-spun ribbons. These precipitates were surrounded by dislocations, confirming their effect in impeding dislocation movement [36]. The precipitation of Mn may be attributed to the uneven element content caused by the addition of Cu. During the sintering process, the precipitated Mn was oxidized to form Mn-O. Although the precipitation was accidental, it had a significant effect on improving the strength of the sintering neck. Due to the lack of stress-induced phase transition properties, the precipitate may have some impact on the superelastic effect, but it significantly improves the strength of the sintered neck. For porous alloys, higher structural strength means that more stress can be transmitted to the grains, thereby preventing premature collapse of the porous alloys due to the fragile sintering neck.

High-resolution transmission electron microscopy (HRTEM) was used to characterize the microstructure of the Cu2 porous alloy shown in Figure 5a, where the matrix consists of austenite and nano-scale martensite phases. The inside diffraction patterns were acquired via the selected area electron diffraction (SAED) technique. The major pattern confirmed that the matrix is the L2₁ austenitic phase. Significantly, there are a set of spots other than the main diffraction point, which indicate the presence of the martensitic phase [14]. In Figure 5a, as well as its enlarged version in Figure 5b, a lath-like area can also be observed, showcasing a typical wavy periodic arrangement of fine substructure. Through the fast Fourier transform (FFT) pattern, seven periodic arrangements of superlattice diffraction spots can be observed between the two basic diffraction spots, indicating the seven-period modulation structure. This confirms the existence of the 7M martensitic structure corresponding to the results obtained by XRD [37]. The existence of the nano-scale martensitic phase can serve as the nucleation point for stress-induced martensitic transformation, which is beneficial for reducing the critical stress of martensitic transformation, thereby improving superelastic performance [38].

Figure 6 shows TEM images of the precipitates in the Cu2 porous alloy at room temperature. Here, the precipitate with a size of around 200 nm was calibrated as MnO. There are obvious protrusions in the dislocation that correspond to the shape of the precipitate. The precipitate embedded within the matrix phase can serve as an impediment in the path of dislocation movement, prompting the dislocation to circumvent the precipitate and proceed with its motion [39]. According to the bright-field image in Figure 6b, there is a precipitate with a size of 5 μm with a clear boundary between the precipitate and matrix. Figure 6c is the HRTEM analysis of the grain boundaries. And Figure 6d,e are the SAED patterns of the surrounding martensite and the precipitate, respectively. The precipitate was MnO with a cubic structure and the spatial group of Fm$\overline{3}$m. It can be seen that the phase boundary is straight and smooth, suggesting a semi-coherent grain boundary between the precipitate and the matrix [40]. It satisfies the semi-coherent orientation relationship of (10$\overline{3}$)_M//($\overline{1}$11)_MnO, (200)_M//($\overline{2}$00)_MnO, and [0$\overline{1}$0]_M//[0$\overline{1}$1]_MnO. Based on the above results, the precipitates formed in porous alloys can impede the movement of dislocations, ultimately enhancing their strength, which aligns with the observations depicted in Figure 4.

For SMAs, high σ_SIMT often results in large energy dissipation and irreversible strain, which leads to a reduction in superelasticity during cyclic loading. At a certain temperature, lower σ_SIMT can greatly improve fatigue resistance, especially for brittle Ni-Mn-based Heusler alloys [41]. Figure 7 illustrates the relationship between the critical stress, critical strain, and preparation state of SMA martensitic transformation in previous studies. For the widely used binary Ni-Ti alloys, the σ_SIMT is generally higher and the critical strain is lower, as illustrated in the ellipse. The Ni-Mn-Ga-Cu specimens prepared by the drop cast method have a higher superelastic strain, shown inside the rectangular box, due to the rapid solidification process that results in a special orientation of the grains. The σ_SIMT of a single crystal and microwire is relatively low, which is due to the lower number of grain boundaries, which reduces the resistance of phase transformation. As the critical strain increases, the critical stress tends to decrease, as shown by the green line. Meanwhile, the σ_SIMT of the porous shape memory alloy prepared in this work is similar to that of a single crystal and microwire (near the green line). This indicates that the incorporation of pores can effectively decrease the number of grain boundaries, subsequently lowering the stress required to induce martensite transformation. Given all this, porous SMAs can be fabricated into large three-dimensional materials to adapt to more scenarios.

In order to analyze the influence of pores on the elastocaloric effect, adiabatic temperature changes (ΔT_ad) of the Cu2 porous and bulk alloys were measured, as shown in Figure 8. During the testing, the Cu2 porous and bulk alloy samples were slowly loaded (0.02 mm/min) and quickly unloaded (5 mm/min), which is considered quasi-adiabatic because the loading time is significantly shorter than the heat convection time. Because of the same loading rate, the loading period of the bulk alloys (500 s) is much longer than that of the porous alloys (100 s). Moreover, the maximum stress of the bulk alloy reaches 700 MPa, which is significantly higher than that of the porous alloys (150 MPa). The ΔT_ad of the Cu2 porous and bulk alloys is shown in Figure 8a,b during 10 loading cycles. The inner graph corresponds to the stress–strain curve. In Figure 8a, the Cu2 porous alloy obtained the ΔT_ad of 0.6 K and maintain stability with a relatively low stress (150 MPa) and strain (1.2%). As a comparison in Figure 8b, the ΔT_ad of the bulk alloy stabilized at 0.8 K, whereas the ΔT_ad of the Ni-Fe-Ga porous alloy reaches 3.4 K [30]. The lower ΔT_ad of the Cu2 porous alloy may be due to asynchronous deformation caused by the lower pore size. Although the ΔT_ad of the Cu2 porous alloy is much lower than that of other excellent elastocaloric materials [55,56], this result is still noteworthy given the low cost due to the low density of the porous alloy. As indicated by the arrow in Figure 8c, the porous alloys exhibit a quicker recovery rate after unloading compared to the bulk alloys, based on the temperature change observed during a single cycle. The above phenomenon indicates that the addition of pores results in a larger surface-to-volume ratio [57], thereby improving the heat exchange efficiency with the environment. In addition, the through-hole structure allows fluid to flow through, which means that the elastocaloric effects of alloys can be more effectively utilized in practical applications.

4. Conclusions

Porous Ni-Mn-Ga SMAs with an average pore size of 20–30 μm were prepared by powder metallurgy using NaCl as a pore space holder. Introducing pores into the Ni-Mn-Ga alloy can effectively reduce the number of grain boundaries in the polycrystalline alloy, thereby decreasing the stress to induce martensite transformation. The addition of Cu forms Mn-O precipitates at the grain boundaries of the sintering neck, blocking the propagation of cracks, thereby enhancing the bonding force between alloy powders and improving the superelasticity. The maximum elongation of the Ni-Mn-Ga porous alloy is 5% and the σ_SIMT is about 160 MPa. This reduction is significant when compared to polycrystalline alloys. And it aligns closely with the characteristics of single crystals and microwires. On the whole, the porous SMAs have a low σ_SIMT, high superelasticity, and relatively low cost, thereby offering advantages in the design of devices for shock absorption. And the porous alloys could become an acceptable alternative to the classical Ti-Ni shape memory alloys in some applications.