Giant Elastocaloric Effect and Improved Cyclic Stability in a Directionally Solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ Alloy

Abstract

Superelastic shape memory alloys with an integration of large elastocaloric response and good cyclability are crucially demanded for the advancement of solid-state elastocaloric cooling technology. In this study, we demonstrate a giant elastocaloric effect with improved cyclic stability in a <001>_A textured polycrystalline (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy developed through directional solidification. It is shown that large adiabatic temperature variation (|ΔT_ad|) values more than 15 K are obtained across the temperature range from 283 K to 373 K. In particular, a giant ΔT_ad up to −27.2 K is achieved by unloading from a relatively low compressive stress of 412 MPa at 303 K. Moreover, persistent |ΔT_ad| values exceeding 8.5 K are sustained for over 12,000 cycles, exhibiting a very low attenuation behavior with a rate of 7.5 × 10⁻⁵ K per cycle. The enhanced elastocaloric properties observed in the present alloy are ascribed to the microstructure texturing as well as the introduction of a secondary phase due to boron alloying.

1. Introduction

Refrigeration systems such as air-conditioners and refrigerators are playing an increasingly important role in modern life. Vapor compression remains the predominant technology in refrigeration facilities, yet the high global-warming potential of volatile liquid refrigerants has raised extensive concerns on environmental degradation [1]. Elastocaloric refrigeration, which is based on the elastocaloric effect (eCE) of solid-state materials, offers an eco-friendly cooling alternative that circumvents potential impacts on the global environment [2,3,4]. Shape-memory alloys (SMAs) show great potential as candidate materials for elastocaloric cooling due to their ability to exhibit an elastocaloric response through the utilization of released and absorbed latent heat in association with stress-induced martensitic transformation. Consequently, very pronounced isothermal entropy change (ΔS_iso) or adiabatic temperature change (ΔT_ad) can be expected upon the application or removal of certain uniaxial stresses [5].

Heusler-type Ni-Mn-X (e.g., X = Ga, In, Sn)-based shape-memory alloys have been shown to exhibit various functional behaviors by utilizing the martensitic transformation, such as magnetic shape memory effect [6,7], magnetocaloric effect [8,9], and elastocaloric effect [10]. However, the cycling life of elastocaloric response in these alloys suffers from strong constraint of poor mechanical properties (e.g., 15 cycles for a Ni₅₀Mn₃₃In_15.5Cu_1.5 alloy [11] and 100 cycles for a Ni_50.4Mn_27.3Ga_22.3 alloy [12]), making them difficult to become reliable elastocaloric refrigerants [10,11,12,13]. The recently developed all-d-metal Heusler-type Ni-Mn-Ti alloys have attracted considerable attention owing to the enhanced mechanical properties that are achieved by the hybridization of d-d atomic orbitals [14]. These alloys exhibit great potential in cooling applications owing to the very large volume change in unit cell (ΔV/V₀) between austenite and martensite, which substantially augments the entropy change ΔS_tr of martensitic transformation and consequently results in very prominent elastocaloric response [15]. In a polycrystalline (Ni₅₀Mn_31.5Ti_18.5)_99.8B_0.2 alloy, it has been reported that removing a compressive loading of 700 MPa can yield a colossal adiabatic temperature drop of up to −31.5 K [15].

Although Ni-Mn-Ti alloys exhibit improved mechanical properties and elastocaloric effect when compared to some other typical Ni-Mn-X-based alloys [10,11,16], there remains a lack of sufficient cyclability in polycrystalline alloys, e.g., fewer than 100 cycles in a Ni₅₀Mn_31.5Ti₁₈Cu_0.5 alloy [17]. The primary cause of this dilemma stems from the constraining effect of adjacent grain boundaries in conventional polycrystalline alloys since the stress-induced martensitic transformation is very sensitive to crystallographic orientation, resulting in deformation incompatibility among grains with different crystallographic orientations. Consequently, significant stress concentration occurs at triple junctions with increasing superelastic cycles, ultimately leading to sample fracture. To enhance the deformation compatibility of neighboring grains during phase transformation, an alternative strategy is to develop a highly textured microstructure with columnar grains through directional solidification [11,18]. Moreover, it is noted that the necessary driving stress for achieving a colossal elastocaloric effect in polycrystalline Ni-Mn-Ti alloys remains too high [14,19]. Since the transformation stress can be effectively manipulated by crystallographic orientation [20,21], microstructure texturing is also expected to reduce the driving stress of martensitic transformation. In addition, alloying new elements is known to be very effective in tailoring the mechanical properties and functional stability [22,23]. Recently, the cyclic number for a (Ni₅₀Mn_30.75Ti_18.25Cu₁)_99.8B_0.2 polycrystalline alloy has been successfully increased to 650 cycles [24], owing to the substantial increase in grain boundary strength due to boron doping. It is known that the precipitate phases can also be introduced by boron doping [25], which could be exploited for the improvements in mechanical properties and functional stability.

To tackle the tough challenges of relatively low cyclability and high driving stress for the elastocaloric effect in polycrystalline Ni-Mn-Ti alloys, a strategy with the combination of microstructure texturing and boron doping was employed in this work. Considering that boron doping may increase the transformation temperatures [15], a Ni₅₀Mn₃₁Ti₁₉ alloy with relatively low transformation temperature (i.e., M_s = 220 K, M_f = 210 K, A_s = 236 K, and A_f = 243 K) was used as the starting alloy to facilitate the exploration of elastocaloric response around room temperature for the boron doped alloys. For the purpose of strengthening the matrix by utilizing the precipitates, 1% boron was introduced and a <001>_A textured polycrystalline (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy with a coarse columnar-grained microstructure was developed by directional solidification. Large |ΔT_ad| values over 15 K were achieved when removing the compressive loading from 283 K to 373 K. Notably, an impressive ΔT_ad of −27.2 K was observed at 303 K upon eliminating a relatively low compressive stress of 412 MPa. Moreover, |ΔT_ad| values exceeding 8.5 K were sustained for over 12,000 cycles, showing improved cyclability owing to the synergistic effects of microstructure texturing and boron doping.

2. Materials and Methods

Polycrystalline (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ (at. %) alloy was fabricated by repeatedly melting high-purity elements in a conventional arc-melting furnace. To obtain a well-oriented microstructure, directional solidification was performed by using the Bridgman method with a temperature gradient of 120 K cm⁻¹ [26], where liquid Ga-In-Sn metal was utilized as the cooling medium. The arc-melted alloy was enveloped in a corundum crucible with an inner diameter of 10 mm and a length of 150 mm, and subsequently remelted at 1563 K. Following a holding time of 1 h, the alloy was directionally solidified by drawing the crucible into a cylinder filled with liquid Ga-In-Sn metal at a constant speed of 3 mm min⁻¹. To eliminate the compositional inhomogeneity, the as-cast alloys were enclosed in vacuum quartz tubes and then subjected to a heat treatment at 1223 K for 48 h. Subsequently, the alloys were rapidly cooled by quenching in cold water. The composition analyses were performed using an energy dispersive spectrometer (EDS). The start and finish temperatures for the forward and reverse martensitic transformation (M_s, M_f, A_s, A_f) were determined through differential scanning calorimetry (DSC-25, TA Instruments, New Castle, USA), using a temperature scanning rate of 10 K min⁻¹. The microstructural features and preferred orientation were characterized using a scanning electron microscope (SEM, JEOL, JSM-IT800, Tokyo, Japan) and the attached electron backscatter diffraction (EBSD) system. To test the mechanical properties, uniaxial compression along the solidification direction (SD) was conducted using a material testing machine with a temperature control unit (AGS-X/50 kN, Shimadzu, Kyoto, Japan). The compressive axis was aligned parallel to the longer edge of a cuboid-shaped sample sized 6 mm × 4 mm × 3 mm. A K-type thermocouple with a resolution of 0.1 K was adhered to the sample surface to track the temperature variation ΔT_ad caused by rapid loading or unloading. The temperature of the sample was continuously recorded with an interval of 0.1 s. A 15 s dwell time was set between loading and unloading to ensure the thermal equilibrium between the sample and the ambience.

3. Results

DSC measurements demonstrate that the martensitic transformation for the directionally solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy occurs below room temperature (Figure 1a) and the corresponding characteristic temperatures are determined to be M_s = 248 K, M_f = 232 K, A_s = 255 K, and A_f = 269 K. In addition, the entropy change ΔS_tr related to martensitic transformation can be evaluated to be 66.7 Jkg⁻¹K⁻¹. Figure 1b plots the compressive stress–strain relation tested at 293 K with a strain rate of 2.8 × 10⁻⁴ s⁻¹. The directionally solidified alloy can withstand a maximum compressive strength of ~1780 MPa, which is much higher than that of the arc-melted alloy (i.e., 1460 MPa, Supplementary Figure S1). In addition, the compressive strength for the present directionally solidified alloy is also higher than those of some other polycrystalline alloys, such as Ni₅₀Mn_31.6Ti_18.4 alloy (i.e., 800 MPa) [27], Ni₅₀Mn_31.75Ti_18.25 alloy (i.e., 1100 MPa) [14], and Ni_36.5Co_13.5Mn₃₅Ti_14.1Gd_0.9 alloy (i.e., 1142 MPa) [28]. Achieving enhanced mechanical properties is of great significance for the functional stability of elastocaloric materials [29].

The global microstructure of the directionally solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy is characterized by large columnar grains, with their longer axes along the SD. Figure 2a displays an EBSD orientation map colored in IPF contrast, covering the region of 5.5 mm × 2.2 mm. It is seen that the columnar grains have a homogeneous orientation with <001>_A along the SD, i.e., <001>_A fiber texture along the SD. As demonstrated in the corresponding inverse pole figure, the texture strength of <001>_A fiber texture is 19.6. The formation of such highly textured microstructure should be attributed to the preferential growth of <001>_A along the temperature gradient direction [30]. In contrast, the arc-melted alloy exhibits equiaxed grains with a very low degree of preferred orientation (Supplementary Figure S2). Thus, it can be inferred that the enhanced mechanical properties of the directionally solidified alloy should be mainly attributed to the columnar-grained microstructure with a strong <001>_A preferred orientation. Such a specific microstructural feature contributes to enhancing the mechanical properties by minimizing stress concentrations at triple junctions and promoting better strain compatibility among adjacent grains. [31]. Microstructural observations also show that certain amounts of precipitated phase appear in the studied alloy, as shown in Figure 2b,c. As determined from the EDS measurements, the matrix has an average composition of Ni_50.3Mn_30.9Ti_18.8, being close to the nominal composition, and the precipitates correspond to a boron-rich phase [32]. Furthermore, a boron-rich secondary phase is observed by TEM, as shown in Figure 2d. It can be observed that the secondary phase is in rod-like shape. The corresponding SAED indicates that the second phase exhibits a face-centered cubic structure (inset of Figure 2d). Similar observations of boron-rich precipitates have also been demonstrated in other boron-doped Ni-Mn-based alloys [13]. The introduction of precipitates has a positive effect in inhibiting crack propagation and alleviating extensive stress concentration [13,24,33,34]. Thus, the enhanced mechanical properties in the present alloy should result from the highly textured microstructure and the introduction of secondary phase due to boron alloying [35,36].

Prior to the characterizations on the superelastic and elastocaloric properties, mechanical training was conducted by performing several cycles of loading–unloading to stabilize stress-induced martensitic transformation (Supplementary Figure S3). Figure 3a presents the superelastic loop for the directionally solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy tested at 293 K. Typical plateau-like superelastic behavior with a stress hysteresis of 105 MPa can be clearly observed. Because of the well-developed <001>_A preferential orientation, a large transformation strain of 4.9% is obtained by stress-induced transformation, enabling a recoverable strain of 8.5% upon unloading.

In order to analyze the elastocaloric properties, the ΔT_ad values obtained by applying/releasing the compressive stress were directly measured. A typical ΔT_ad profile is displayed in Figure 3b, showing an adiabatic temperature rise on loading and a temperature drop upon unloading. Figure 3c depicts the variation in ΔT_ad values with compressive strains detected under various strain rates. The |ΔT_ad| values increase with increasing compressive strain owing to the gradually enhanced proportion of transformed martensite. A rapid increase in |ΔT_ad| values is presented until the compressive strain of 7% and then the increasing rate becomes slow. When applying a compressive strain of 8%, the ΔT_ad values resulting from the application and release of compressive stress at a strain rate of 2.2 × 10⁻² s⁻¹ reach 17.8 K and −17.2 K, respectively. Figure 3d presents the correlation between ΔT_ad and strain rate under various compressive strains. Apparently, increasing the strain rate results in a gradual rise in |ΔT_ad| values on account of the improved adiabatic condition. Thus, both the applied strain and strain rate strongly affect the stress-induced ΔT_ad values. The corresponding stress–strain curves for ΔT_ad measurements under various compressive strains and strain rates are displayed in Figure 4a.

It is noted that applying a higher compressive stain and strain rate also widens the stress hysteresis, resulting in enhanced energy dissipation (ΔW) that is represented by the enclosed region of the superelastic loop, as demonstrated in Figure 4. Such energy dissipation is irreversible, thereby weakening the energy conversion efficiency [37]. An important index used for evaluating the energy efficiency is the coefficient of performance of material (COP_mat), which can be calculated by Equation (1) [38].

$${\mathit{COP}}_{\mathrm{mat}}=|\Delta Q/\Delta W|=|(\rho \Delta T_{\mathit{ad}}·C_p)/\oint \sigma \mathrm{d}\epsilon |$$

(1)

where ρ denotes the density (ρ = 7278 kg m⁻³) and C_p is the specific heat capacity (C_p = 423 Jkg⁻¹K⁻¹, Supplementary Figure S4). Based on the ΔT_ad and ΔW mentioned above, the influence of strain level and strain rate on the COP_mat was analyzed, as summarized in Figure 5. Generally, the COP_mat values are gradually increased with the strain rate at a constant compressive strain. However, when the strain rate is fixed, the COP_mat values are significantly reduced when increasing the compressive strain, even though there is a significant improvement in ΔT_ad.

In order to examine the influence of testing temperature on the superelastic response, compressive tests were conducted over a temperature range from 273 K to 383 K, by using a compressive strain of 8% at a strain rate of 2.8 × 10⁻⁴, as shown in Figure 6a. At each testing temperature, the alloy demonstrates a typical plateau-type superelasticity and a reversible strain of 8% can be clearly observed. With the increase in testing temperature, a higher level of driving stress is needed to activate the martensitic transformation owing to the increased stability of austenite. Consequently, the critical stress σ_cr shows a linearly increasing trend, with a rate of 6.0 MPa K⁻¹ (inset of Figure 6a). Similar linear dependence has also been reported in some other Ni-Mn-Ti alloys [15,39]. According to the stress–strain curves obtained at different temperatures, the compressive loading-induced ΔS_iso was evaluated through the Maxwell relation (Equation (2)).

$$\Delta S_{\mathit{iso}}={ v}_0{\int }_0^{\epsilon }{\left({\displaystyle \frac{\partial \sigma }{\partial T}}\right)}_{\epsilon }d\epsilon$$

(2)

In the equation, v₀ refers to the specific volume (i.e., 1.38 × 10⁻⁴ m³kg⁻³), as shown in Figure 6b. Under the compressive strain of 8%, a maximum ΔS_iso of 38.8 Jkg⁻¹K⁻¹ can be achieved at 303 K.

The unloading ΔT_ad profiles measured at the temperature span from 273 K to 383 K when removing a compressive strain of 8% are presented in Figure 6c. Here, the unloading process was performed by using a much higher strain rate of 0.3 s⁻¹, thereby allowing the enhanced adiabatic condition. A giant elastocaloric response, characterized by |ΔT_ad| values exceeding 15 K, was obtained within the testing temperature region from 283 K to 373 K. It is worth noting that the most remarkable elastocaloric response appears at 303 K, showing the maximum ΔT_ad up to −27.2 K. Such ΔT_ad is in good agreement with the value evaluated from the ΔS_iso (i.e., −27.6 K) based on Equation (3).

$$\Delta T_{\mathit{ad}}^{\mathit{cal}}=\Delta S_{\mathit{iso}}·T_{\mathit{test}}/C_p$$

(3)

Table 1. Comparison of the elastocaloric properties between the present alloy and some other elastocaloric materials.

Alloy	Sample Status	Ttest (K)	|ΔTad| (K)	σmax (MPa)	Strain Rate (s−1)	Ref.
(Ni50Mn31Ti19)99B1	Polycrystal (textured)	303	27.2	412	3.0 × 10−1	This work
Ni50Mn31.75Ti18.25	Polycrystal (textured)	293	20.4	913	2.0	[14]
Ni50Mn35Ti15	Polycrystal (untextured)	293	4.5	450	2.8 × 10−2	[19]
Ni50Mn31.6Ti18.4	Single crystal	293	29.0	830	2.0	[27]
Ni37Co9Fe4Mn35Ti15	Polycrystal (untextured)	293	6.3	400	1.4 × 10−1	[57]
(Ni50Mn31.5Ti18.5)99.8B0.2	Polycrystal (untextured)	308	31.5	700	5.33	[15]
Ni55Mn18Ga27	Polycrystal (textured)	300	10.7	350	2.0 × 10−1	[40]
Ni50Mn18.5Ga25Cu6.5	Polycrystal (textured)	315	8.1	100	4.2 × 10−2	[41]
Ni50(Mn28.5Cu4.5)(In14Ga3)	Polycrystal (textured)	293	19.0	710	2.0	[42]
Ni44Mn46Sn10	Polycrystal (textured)	320	18.0	390	3.0 × 10−1	[10]
Ni45Mn44Sn11	Polycrystal (textured)	298	10.0	310	3.0 × 10−2	[43]
Ni47.5Co4.2Mn37.3Sb12.8	Polycrystal (textured)	303	8.7	400	5.0 × 10−2	[44]
Cu71.3Al17.5Mn11.2	Single crystal	293	11.9	120	1.4	[45]
Cu59.1Zn27Al13.8Zr0.1	Single crystal	343	14.2	200	2.0 × 10−1	[46]
Ni48.9Ti51.1	Wire	322	21.0	900	2.0 × 10−1	[47]
Ni50.8Ti49.2	Polycrystal (textured)	323	17.9	913	2.0 × 10−1	[46]

compares the elastocaloric properties between the present alloy and some other elastocaloric materials. It can be seen that the |ΔT_ad| value (i.e., 27.2 K) of the current alloy is higher than those of Ni-Mn-X alloys (X = Ga, In, Sn and Sb) [10,40,41,42,43,44], Cu-based alloys [45,46], and Ni-Ti alloys [46,47], but lower than those of certain Ni-Mn-Ti alloys. It should be noted that the colossal elastocaloric effect in those Ni-Mn-Ti alloys requires a quite large driving stress, indicating a high level of energy consumption. By comparison, the required external stress for the present alloy is substantially decreased, i.e., 412 MPa, thereby favoring the miniaturization and energy-saving of refrigeration equipment [48]. To quantify the elastocaloric response, the adiabatic temperature change under unit stress was determined by |ΔT_ad/σ_max|. Notably, the maximum |ΔT_ad/σ_max| of the present alloy reaches 78.8 K GPa⁻¹, owing to a considerably large ΔT_ad of −24.9 K realized at a relatively low compressive stress of 316 MPa at 293 K. By comparison, the present |T_ad/σ_max| is also larger than those of some other elastocaloric materials [3,10,11,15,16,20,31,39,40,49,50,51,52,53,54,55,56], as illustrated in Figure 6d.

Figure 6. (a) Temperature-dependent superelasticity of the present alloy subjected to a strain rate of 2.8 × 10⁻⁴ s⁻¹ by applying a maximum compressive strain of 8%. The inset displays the evolution of critical stress σ_cr as the variation of temperature. (b) ΔS_iso induced by compressive loading under various compressive strains. (c) ΔT_ad profiles upon unloading measured at various temperatures. (d) Comparison of the elastocaloric properties between the present alloy and several typical elastocaloric materials [3,10,11,15,16,20,31,39,40,49,50,51,52,53,54,55,56].

Figure 6. (a) Temperature-dependent superelasticity of the present alloy subjected to a strain rate of 2.8 × 10⁻⁴ s⁻¹ by applying a maximum compressive strain of 8%. The inset displays the evolution of critical stress σ_cr as the variation of temperature. (b) ΔS_iso induced by compressive loading under various compressive strains. (c) ΔT_ad profiles upon unloading measured at various temperatures. (d) Comparison of the elastocaloric properties between the present alloy and several typical elastocaloric materials [3,10,11,15,16,20,31,39,40,49,50,51,52,53,54,55,56].

To investigate the cyclability of eCE for the present alloy, cyclic tests under a compressive strain of 4% were performed by using a strain rate of 1.1 × 10⁻² s⁻¹. Figure 7a displays the evolution of stress–strain correlation during the cyclic tests. The superelastic response demonstrates a good stability, with a small stress drop of 19 MPa over 12,000 cycles of testing. Such a stress drop could be caused by the accumulation of transformation-induced defects that stabilize the martensite phase [55,58]. It is noted that there is no indication of fracture on the sample after 12,000 cycles, which is also a reflection of enhanced mechanical properties in the present directionally solidified alloys. Figure 7b presents the representative ΔT_ad profiles corresponding to the various selected cycles. The present alloy exhibits persistent |ΔT_ad| values larger than 8.5 K for more than 12,000 cycles. The degradation rate of |ΔT_ad| values during such long-term tests is very low, i.e., 7.5 × 10⁻⁵ K per cycle. For a purpose of comparative analyses, the cyclic tests were also conducted in a directionally solidified Ni₅₀Mn₃₁Ti₁₉ alloy without boron doping under the same testing parameters. As shown in Figure S5a, the stress–strain curves demonstrate a significant stress drop of 167 MPa after 40 mechanical cycles for the sample. Additionally, the |ΔT_ad| values during unloading also exhibit a sharp decrease from 8.4 K to 4.1 K (Supplementary Figure S5b). Thus, boron doping indeed yields a prominent cyclability of the elastocaloric response. As summarized in

Table 2. Cyclability and degradation rate of the elastocaloric effect for the present alloy and some elastocaloric materials.

Alloy	Sample Status	Number of Cycles	Degradation Rate of ΔTad (K per Cycle)	Ref.
(Ni50Mn31Ti19)99B1	Polycrystal (textured)	12,000	7.5 × 10−5	This work
Ni50Mn30Ti20	Polycrystal (textured)	2000	5.0 × 10−4	[39]
Ni37Co9Fe4Mn35Ti15	Polycrystal (untextured)	1000	2.0 × 10−4	[57]
Ni50.4Mn27.3Ga22.3	Polycrystal (textured)	250	1.2 × 10−3	[66]
Ni50.4Mn27.3Ga22.3	Polycrystal (textured)	100	2.0 × 10−3	[12]
Ni50Fe19Ga27Co4	Single crystal	3000	9.3 × 10−4	[20]
Ni54Fe19Ga27	Single crystal	100	1.0 × 10−3	[67]
Ni54Fe19Ga27	Polycrystal (untextured)	100	9.0 × 10−3	[59]
Ni53.2Fe19.4Ga27.4	Polycrystal (untextured)	480	2.3 × 10−3	[60]
Cu71.3Al17.5Mn11.2	Single crystal	50	4.0 × 10−3	[45]
Cu71Al18Mn11	Polycrystal (textured)	275	1.5 × 10−3	[61]
Cu71.1Al17.2Mn11.7	Polycrystal (textured)	200	1.0 × 10−3	[53]
Cu59.1Zn27Al13.8	Polycrystal (untextured)	10,000	1.4 × 10−4	[46]
Co49Fe3V33Ga15	Polycrystal (textured)	200	1.5 × 10−3	[62]
Co50V35Ga14Ni1	Polycrystal (untextured)	4000	8.5 × 10−4	[63]
(Ni42.5Ti50Cu7.5)99Co1	Polycrystal (textured)	200	2.2 × 10−2	[64]
Ti54.9Ni32.5Cu12.6	Polycrystal (untextured)	1502	2.0 × 10−4	[65]

, both the cycle number and degradation rate for the present alloy are substantially improved when compared to some typical elastocaloric materials reported in the literature [12,20,39,45,46,53,57,59,60,61,62,63,64,65,66,67], thereby showing an enhanced functional stability of the elastocaloric response.

To further understand the degradation of functional behaviors, the DSC curves between the initial sample and the sample after mechanical cycles were compared, as shown in Figure 8. It is seen that mechanical cycling leads to a broadening of the endothermic/exothermic peak, which could be an indication of the accumulation of dislocations and the resulting internal stress [68]. In addition, ΔS_tr is also reduced after long-term testing, i.e., from 66.7 Jkg⁻¹K⁻¹ for the initial sample to 59.8 Jkg⁻¹K⁻¹ in the cycled sample, suggesting the generation of residual martensite induced by mechanical cycling.

Figure 9a shows the microstructure image for the longitudinal section of the directionally solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy after 12,000 mechanical cycles. It is noted that some longitudinal cracks have nucleated and propagated along the loading direction. The occurrence of these cracks can be attributed to the generation of stress concentration at the grain boundary during mechanical cycling. With increasing the cycles, cracks nucleate at the grain boundary and subsequently propagate into the interior of the grains, as highlighted in red dashed squares in Figure 9a. It is worth mentioning that residual martensite can also observed in the alloy, as shown in Figure 9b, which should account for the reduced ΔS_tr observed in the DSC measurements after mechanical cycling. Accordingly, the reduction in the volume fraction of transforming austenite involved in the phase transition results in the attenuation of ΔT_ad during mechanical cycling.

To elucidate the underlying fatigue mechanisms behind long-term cyclic testing, the morphological characteristics of the fractured sample were analyzed. Figure 9c displays the morphology of the longitudinal section for the directionally solidified (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy after 13,000 mechanical cycles. The abundance of river patterns and cleavage steps suggests that the sample exhibits a transgranular fracture behavior, which aligns with the observed crack propagation observed in Figure 9a. Moreover, the cross-section exhibits numerous white granular precipitates that are dispersedly distributed along the grain boundary and within the grain. This provides clear evidence of enhanced cyclic stability by boron doping. These precipitates effectively mitigate deviatoric stress through plastic deformation, thereby impeding crack initiation and propagation [34].

The results presented above indicate that the present directionally solidified alloy exhibits a giant elastocaloric response and improved cyclic stability, which should be attributed to the effective microstructure tuning through directional solidification and boron doping. It is known that during mechanical cycling, dislocations and microcracks tend to preferentially form at fragile grain boundaries, resulting in premature intergranular fracture [69]. By utilizing directional solidification, a coarse-grained microstructure with <001>_A preferential orientation formed in the present alloy significantly reduces the amount of grain boundaries, thereby minimizing crack initiation sites and enabling prolonged service life [70,71]. Moreover, the well-oriented microstructure is also favorable for the improvement in the deformation compatibility during stress-induced phase transformation, which thus weakens the grain boundary constraints on the lattice deformation and contributes to a giant elastocaloric response under reduced driving stress [11]. On the other hand, boron is recognized as an electron donor [72,73], and its microalloying can effectively enhance the cohesive strength of grain boundaries, thereby increasing the mechanical properties by impeding the intergranular fracture. This effect has been observed in various structural intermetallic compounds such as Ni₃Al [74], Ni₃Fe [75], and FeAl [76] alloys. Here, a boron-rich secondary phase is also introduced in the present alloy due to boron doping. This secondary phase is distributed within the grain and along the grain boundary (Figure 2c), which is favorable for enhancing the mechanical properties and thus extending the service life. Overall, the enhanced elastocaloric response for the present directionally solidified alloy should primarily arise from the synergy of microstructure texturing and boron doping.

4. Conclusions

In summary, the elastocaloric effect and its functional stability in a <001>_A textured polycrystalline (Ni₅₀Mn₃₁Ti₁₉)₉₉B₁ alloy fabricated by directional solidification were investigated. Through boron doping and microstructure control, the mechanical properties were significantly enhanced, giving rise to the maximum compressive strength of 1780 MPa. Consequently, a giant elastocaloric response represented by |ΔT_ad| values higher than 15 K was realized over a wide temperature window ranging from 283 K to 373 K. Because of the reduced driving stress derived from the highly textured microstructure, a very high value of specific adiabatic temperature change (|ΔT_ad/σ_max|) up to 78.8 K GPa⁻¹ was achieved. Moreover, persistent |ΔT_ad| values exceeding 8.5 K were sustained for over 12,000 cycles, showing enhanced cyclability when compared to some other elastocaloric materials. It is demonstrated that the combination of microstructure texturing and boron doping can be exploited as an effective approach in enhancing the cyclic stability of the elastocaloric effect.