Introduction

Refrigeration is a cornerstone of modern infrastructure, enabling technologies from food preservation to electronics and medical systems1. However, conventional vapor compression systems are nearing their theoretical efficiency limits and rely on refrigerants with high global warming potential, raising serious environmental concerns2,3. These challenges have spurred the development of sustainable, high-efficiency solid-state cooling technologies. Among these, elastocaloric cooling, which leverages the reversible martensitic transformation in shape memory alloys (SMAs), has emerged as a particularly promising solution due to its high energy efficiency and environmentally benign nature4,5,6,7,8.

The elastocaloric effect is driven by the latent heat exchange associated with stress-induced martensitic transformation9, with performance typically quantified by the adiabatic temperature change \(\left({\Delta T}_{{ad}}\right)\) 10,11. TiNi-based SMAs are among the most extensively studied due to their excellent mechanical and functional properties12,13. However, real-world elastocaloric applications require these materials to maintain functionality over millions of mechanical cycles14. TiNi alloys, in particular, are susceptible to both structural fatigue—leading to fracture—and functional fatigue, characterized by a decline in elastocaloric performance over time15,16. Thus, achieving both a large elastocaloric response and long-term cyclic durability remains a fundamental challenge.

A variety of strategies have been explored to address this issue, including compositional tuning, grain refinement, precipitation hardening, and composite architectures17,18,19,20,21,22,23,24. Among these, precipitation engineering has shown particular promise21,22,25. Strengthening phases such as Ti2Cu17, Ti2Ni26,27, Ti2Ni322, and TiNi321 have been introduced to enhance fatigue life by promoting transformation reversibility at matrix–precipitate interfaces. However, excessive precipitation can hinder martensitic transformation, thereby diminishing the elastocaloric response21,22. Thus, optimizing the balance between mechanical strengthening and transformation compatibility is essential.

Crystallographic orientation also plays a critical role in dictating the functional response of TiNi alloys. Shape memory and superelastic behaviors are both strongly orientation-dependent28,29, and this extends to elastocaloric performance. Techniques such as cold rolling and directional solidification have been used to induce strong texture, with significant variation in elastocaloric response observed along different loading directions30,31,32. Recent reports of high elastocaloric performance in highly textured Heusler alloys and Ti-based SMAs—including NiMnTi and Ti-Nb-Zr-Ta—as well as enhanced stability in tweed-textured Fe-Pd single crystals, highlight the power of texture control33,34,35,36,37,38,39,40. Yet, limited attention has been paid to integrating crystallographic texture with microstructural engineering to improve fatigue resistance in elastocaloric TiNi alloys.

Notably, TiNi alloys exhibit strong mechanical asymmetry: while <001 > B2-oriented grains show limited transformation under tension, they exhibit large recoverable strains under compression41,42. Moreover, introducing coherent precipitates with specific orientation relationships has been shown to promote reversible martensitic transformation and suppress fatigue damage by enabling martensite nucleation via localized lattice distortions17,21,22. These insights suggest that co-designing texture and precipitate architecture could be a powerful strategy to simultaneously achieve high elastocaloric performance and exceptional fatigue resistance.

Here, we report a textured Ti49Ni₅₁ (at. %) alloy with high-density, epitaxially aligned Ti₄Ni₂O precipitates, synthesized via controlled directional solidification. The alloy exhibits a large \({\Delta T}_{{ad}}\) of −15.9 K, sustained over 10⁷ compressive loading-unloading cycles. This performance is proposed to derive from synergistic effects of strong crystallographic texture and a distributed martensitic transformation enabled by interfacial lattice distortions at B2/Ti₄Ni₂O boundaries. These findings underscore the promise of texture–precipitate co-design in TiNi-based SMAs and offer a viable pathway for developing durable, high-performance elastocaloric materials for solid-state cooling.

Results

Textured microstructure with aligned Ti₄Ni₂O precipitates enables large and stable transformation strain

The microstructure of the directionally solidified Ti49Ni51 alloy was first characterized using scanning electron microscopy (SEM). Backscattered electron (BSE) images captured along and perpendicular to the solidification direction (SD) reveal a highly textured structure comprising columnar grains oriented along the SD (Fig. 1a, b). The grain lengths exceed millimeter with widths ranging from tens to hundreds of micrometers. Strong contrast variation between grains in the cross-sectional view (perpendicular to SD) indicates substantial orientation differences, whereas the longitudinal section (parallel to SD) exhibits minimal contrast, suggesting a preferred grain orientation along the SD. Notably, fine precipitates are uniformly distributed throughout the columnar grains, aligned along the SD, and without significant segregation near grain boundaries.

Fig. 1: Microstructural characterization of the as-prepared Ti49Ni51 alloy.
Fig. 1: Microstructural characterization of the as-prepared Ti49Ni51 alloy.
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BSE SEM images acquired parallel (a) and perpendicular (b) to the SD, revealing a columnar grain structure aligned with SD. c XRD pattern confirms the coexistence of the B2 matrix and Ti₄Ni₂O precipitates. d Bright-field TEM image with corresponding EDS elemental maps, highlighting the composition of the Ti₄Ni₂O precipitate embedded within the B2 matrix. e EDS line scan along the red arrow in (d), showing the elemental distribution and interface sharpness across the matrix–precipitate boundary. Source data are provided as a Source Data file.

X-ray diffraction (XRD) confirms the presence of the B2 austenite phase along with a secondary phase indexed as Ti₄Ni₂O (Fig. 1c). No martensitic phase is detected at room temperature. The Ti₄Ni₂O phase adopts a cubic structure (space group \({{{\rm{Fd}}}}\bar{3}{{{\rm{m}}}}\)), similar to the Ti₂Ni phase, with only slight differences in lattice parameters43,44. To further identify the chemical composition of the precipitates, scanning/transmission electron microscopy (S/TEM) coupled with energy-dispersive X-ray spectroscopy (EDS) was performed. A representative bright-field TEM image (Fig. 1d) shows a submicron precipitate embedded in the matrix. Elemental mapping and line profile analysis across the interface (Fig. 1e) reveal that the precipitate is enriched in Ti and O but depleted in Ni, with an approximate atomic ratio of Ti:Ni:O = 4:2:1. This confirms the formation of Ti₄Ni₂O precipitates coherently embedded within the B2 matrix.

Electron backscatter diffraction (EBSD) analysis performed on the transverse cross-section (perpendicular to SD) reveals a strong crystallographic texture in both the B2 matrix and the Ti₄Ni₂O precipitates. The inverse pole figure (IPF) map (Fig. 2a) indicates a predominant orientation close to <001 > B2 aligned with the SD. Several discrete submicron regions are identified as Ti₄Ni₂O precipitates (Supplementary Fig. S1), which also share the <001> orientation. Pole figure (PF) analysis confirms an identical orientation relationship between the matrix and precipitates: {001} < 110 > B2 // {001} < 110>Ti₄Ni₂O (Figure S2). This mutual texture supports the epitaxial growth of Ti₄Ni₂O precipitates within the preferentially oriented B2 grains during directional solidification.

Fig. 2: Textured microstructure and orientation-dependent mechanical behavior.
Fig. 2: Textured microstructure and orientation-dependent mechanical behavior.
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a IPF orientation map of the B2 phase acquired on a cross-section perpendicular to the SD, revealing a strong texture along the SD. The corresponding IPF color legend is shown in the inset. b Theoretical transformation strain map under compression along SD, derived from the orientation data in (a) based on the crystallographic relationship between the B2 parent phase and B19′ martensite. The average calculated transformation strain reaches 6.1%. c Ambient temperature compressive stress–strain curves measured parallel and perpendicular to SD, highlighting significant anisotropy in mechanical response and enhanced superelasticity along SD. Source data are provided as a Source Data file.

Given the orientation-dependent transformation behavior in TiNi alloys, the local transformation strain was analyzed based on the measured orientation distribution and the known crystallography of the B2 to B19′ martensitic transformation12. A two-dimensional transformation strain map derived from the grain orientation data (Fig. 2b, Supplementary Fig. S3) shows an average theoretical transformation strain of ~6.1% under compression along the SD, with minor variations across different grains due to slight orientation deviations. Theoretical averaged transformation strains for arbitrary loading directions were computed from EBSD data and indicate a clear orientation preference (Figure S3).

To assess the mechanical response, superelastic cycling tests were conducted at room temperature in both parallel and perpendicular directions relative to the SD. Samples tested along the SD exhibit a large recoverable strain of up to 6.2%, maintaining stable superelasticity over repeated cycles (Fig. 2c). In contrast, samples tested perpendicular to the SD accumulate significant residual strain within 20 cycles (Fig. 2c), indicating poor functional stability. Comparing experimental results with these theoretical averages in Fig. S3, the measured compressive strain along the SD (6.2%) corresponds to ~102% of the theoretical average (6.1%), with the slight excess (~0.1%) attributable to elastic strain under high applied stress. In contrast, the strain perpendicular to the SD decreases from 4.5% initially to a stabilized value of 3.2% after 20 cycles, corresponding to ~80% and ~57% of the theoretical average (5.6%), respectively. The parallel direction also shows a nonlinear, narrow hysteresis superelastic response without a pronounced stress plateau, accompanied by a reduced elastic modulus (~38.3 GPa) compared to the perpendicular direction (~63.2 GPa). This reduced modulus and diminished hysteresis suggest a regulated and anisotropic martensitic transformation under the texture-precipitate synergy. These results highlight the superior cyclic stability and recoverability of the textured direction, which is critical for long-term elastocaloric performance.

Elastocaloric effect with high fatigue life

The elastocaloric performance and fatigue resistance of the textured Ti49Ni51 alloy were evaluated via cyclic compressive loading along the textured direction (parallel to the SD) at room temperature. As shown in Fig. 3a, b, the alloy exhibits exceptional durability, enduring up to 10⁷ compressive cycles without fracture. Despite minor attenuation of the superelastic response over cycling, the stress–strain curves (Fig. 3a) show a substantial recoverable strain of 5.1% after 10⁷ cycles, with total residual strain of ~0.75%. Both the hysteresis and elastic modulus decrease gradually during cycling, suggesting evolving but stable transformation behavior. A crossover emerges at ~4% strain in the stress–strain curves after ~10⁶ cycles, originating from the onset of a two-step transformation (Figure S11). The \({\Delta T}_{{ad}}\), a direct measure of the elastocaloric effect, was also monitored throughout the fatigue life (Fig. 3b). Initially, the alloy delivers a large \({\Delta T}_{{ad}}\) of −21.8 K. After an initial drop of 5.7 K within the first 2 × 10⁶ cycles, \({\Delta T}_{{ad}}\) stabilizes, decreasing by only 0.2 K over the remaining 8 × 10⁶ cycles. This stabilization indicates that most performance degradation occurs early in cycling, after which the elastocaloric response remains largely constant, highlighting excellent long-term functional reliability.

Fig. 3: Fatigue endurance and elastocaloric performance of the textured Ti49Ni51 alloy under compression along the SD.
Fig. 3: Fatigue endurance and elastocaloric performance of the textured Ti49Ni51 alloy under compression along the SD.
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a Superelastic stress–strain curves obtained after various numbers of compressive fatigue cycles at a maximum stress of 900 MPa, demonstrating excellent strain recoverability. b Evolution of the \({\Delta T}_{{ad}}\) after the same number of cycles as in (a), recorded under rapid unloading from 900 MPa, indicating stable elastocaloric response. c, d Comparison of the fatigue performance of \({\Delta T}_{{ad}}\) and recoverable strain (εrec) in the textured Ti49Ni51 alloy with reported bulk SMAs, including nanostructured TiNi₅₀.₈ alloy19, TiNi/TiNi3 eutectic TiNi alloy21, hypoeutectic TiNi58 alloy22, nanocrystalline Ti-44Ni-5Cu-1Al alloy23, TiNi50.8 spring45, polycrystalline Ni50Ti45.3V4.7 alloy46, pre-strained TiNi50.9 plates47, Cu and B co-doped Co-Ni-Ga alloy48, highlighting the superior fatigue resistance and elastocaloric stability of the present work. Source data are provided as a Source Data file.

A comparative analysis of bulk SMAs reported in literature19,21,22,23,45,46,47,48 (Fig. 3c) underscores the exceptional balance achieved in this textured Ti49Ni51 alloy. Unlike typical SMAs that display a trade-off between elastocaloric strength and fatigue life, the present material simultaneously exhibits an ultrahigh fatigue life (10⁷ cycles), a substantial elastocaloric effect (\({\Delta T}_{{ad}}\) = −15.9 K), and a large coefficient of material (COPmat = 23.7 after fatigue, Figure S5). Moreover, a relatively low compressive driving stress of ≈900 MPa is sufficient to induce complete transformation in our textured Ti₄₉Ni₅₁ alloy compared with several strengthened TiNi systems19,21,22, reducing mechanical energy input and enhancing potential for practical applications, although further optimization is required for commercialization. Importantly, the recoverable strain remains significant even after extended cycling (Fig. 3d), further supporting the material’s robust functional endurance. These combined attributes—large and stable \({\Delta T}_{{ad}}\), long fatigue life, high COPmat—establish the textured Ti49Ni51 alloy as a highly promising candidate for durable elastocaloric cooling applications.

Uniform and progressive martensitic transformation upon loading

To elucidate the origin of the ultrahigh fatigue resistance, the phase transformation behavior during loading was characterized in detail. Digital image correlation (DIC) mapping of strain fields (Fig. 4a) reveals a transformation strain distribution of over 6% under compression. While local strain variations are present, the overall transformation field is more spatially uniform compared with the pronounced strain localization and Lüders-like banding during stress-induced martensitic transformation (SIMT) typically observed in conventional TiNi alloys49,50. This observation aligns with the previously reported nonlinear stress–strain response lacking a distinct plateau (Fig. 2), consistent with macroscopically uniform transformation phenomena observed in select TiNi alloys51,52.

Fig. 4: Uniform and progressive martensitic transformation under compressive loading in the textured Ti49Ni51 alloy.
Fig. 4: Uniform and progressive martensitic transformation under compressive loading in the textured Ti49Ni51 alloy.
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a DIC strain maps captured during compressive loading and unloading up to 900 MPa reveal spatially uniform transformation strain distribution, indicating a homogeneous transformation process. b In-situ XRD patterns acquired under stepwise compressive loading and unloading show a gradual evolution of diffraction peaks corresponding to the B2 and B19′ phases. c Relative diffraction intensities of the (100)B2 and (100)B19′ reflections are extracted from peak area analyses, illustrating a smooth and progressive transformation between B2 and B19′ phases during the loading-unloading cycle. The relative diffraction intensity of (100)B2 and (100)B19′ are defined as \(\frac{{I\left(100\right)}_{B2}}{{I(100)}_{B2}+{I(100)}_{B19{\prime} }}\times 100\%\) and \(\frac{{I(100)}_{B19{\prime} }}{{I(100)}_{B2}+{I(100)}_{B19{\prime} }}\times 100\%\), respectively. The error bars represent the standard error obtained from peak analysis. Source data are provided as a Source Data file.

To further probe the transformation pathway, in-situ XRD analysis was performed under incremental compressive stress up to 900 MPa (Fig. 4b). Initially, diffraction peaks correspond exclusively to the B2 parent phase and Ti₄Ni₂O precipitates. As stress increases, peaks associated with B19′ martensite gradually appear and intensify, while those of the B2 phase diminish. Upon unloading, the diffraction profile nearly returns to its original state, confirming a highly reversible martensitic transformation with minimal retained martensite. The evolution of the (100)B2 and (100)B19’ peak intensities is plotted in Fig. 4c. The relative intensity of (100)B19’ increases smoothly with stress beyond 500 MPa, reaching a maximum at peak load, and then gradually diminishes during unloading. This gradual and continuous evolution confirms the absence of abrupt martensite nucleation or burst-like propagation, distinguishing the transformation from classical SIMT and indicating a diffuse, progressive transformation mechanism. To avoid confusion with the strict thermodynamic concept of a second-order transition, we use ‘progressive’ or ‘distributed’ transformation to denote a macroscopically smooth, multi-site B2 → B19′ transformation that proceeds over a wide stress interval while retaining two-phase coexistence at microscopic length scales. Such a progressive phase transformation under wide stress range minimizes internal stress concentrations and reduces microstructural damage, thereby enhancing transformation reversibility and fatigue life. The progressive martensite formation and reversion also help suppress defect accumulation that typically degrades performance in conventional SMAs. These insights underscore the pivotal role of microstructure-controlled transformation pathways in achieving superior fatigue resistance, warranting further investigation into the nanoscale characteristics governing this behavior.

Promoted nucleation and confined growth of martensite under lattice distortions near B2/Ti4Ni2O interface

To clarify the underlying microstructural mechanism responsible for the observed transformation behavior, in-situ TEM observations were performed during cooling from 315 K to 105 K. As shown in Fig. 5a, the microstructure at 315 K—well above the martensitic transformation temperature—consists of a B2 matrix with Ti₄Ni₂O precipitates. The selected-area electron diffraction (SAED) pattern (Fig. 5a′) from the red-circled region confirms the presence of a single-phase B2 matrix between precipitates at this temperature. Upon cooling to 215 K, just below the martensitic transformation start temperature, nucleation of B19′ martensite is observed in proximity to B2/Ti₄Ni₂O interfaces (Fig. 5b). The corresponding SAED pattern (Fig. 5b′) reveals the coexistence of both B2 and B19′ phases in this region, suggesting that the interface serves as a favorable nucleation site. At 105 K—below the martensitic finish temperature—the matrix between precipitates is nearly fully transformed into martensite (Fig. 5c). The SAED pattern (Fig. 5c′) confirms the presence of a single-variant B19′ martensitic structure, devoid of twinning. These observations suggest that the B2/Ti₄Ni₂O interfaces facilitate the directional nucleation and confined growth of martensite, enabling a smooth and spatially uniform transformation throughout the microstructure. The absence of twinned martensite further supports the hypothesis of interface-mediated variant selection, likely governed by strain fields at the phase boundaries.

Fig. 5: In-situ TEM observation of B19′ martensite nucleation and growth near the B2/Ti₄Ni₂O interface during cooling.
Fig. 5: In-situ TEM observation of B19′ martensite nucleation and growth near the B2/Ti₄Ni₂O interface during cooling.
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Bright-field TEM images acquired at progressively decreasing temperatures of 315 K (a), 215 K (b), and 105 K (c) show the evolution of martensitic transformation. A nano-scale thin layer of B19′ martensite begins to nucleate near the B2/Ti₄Ni₂O interface at 215 K (b), with further growth observed at 105 K (c). a′–c′ Corresponding SAED patterns taken from the region marked by red circles in (a–c). Additional diffraction spots associated with the emergent B19′ martensite phase appear in (b′) and are highlighted by red arrows, confirming the phase transformation onset.

To investigate the structural features governing this behavior, high-resolution transmission electron microscopy (HR-TEM) was employed to examine the B2/Ti₄Ni₂O interface (Fig. 6). The bright-field image (Fig. 6a) and corresponding SAED pattern (Fig. 6a′) confirm a well-defined orientation relationship of {100} < 011 > B2 // {400} < 011>Ti₄Ni₂O between the matrix and precipitate. A closer examination of the B2/Ti4Ni2O interface region (red square in Fig. 6a) is shown in Fig. 6b, accompanied by its Fast Fourier Transform (FFT, Fig. 6c), where the presence of streaking indicates substantial lattice distortion near the interface. Further analysis via inverse FFT (IFFT), generated by selecting the overlapping (100)B2 and (400)Ti₄Ni₂O reflections (red circle in Fig. 6c) is shown in Fig. 6d, revealing a high density of dislocations localized at the interface. The lattice strain distribution near the B2/Ti₄Ni₂O boundary was quantitatively mapped using geometric phase analysis (GPA)53, shown in Fig. 6e–g. Strain mapping reveals localized lattice distortions exceeding ±10% adjacent to the interface, confirming the presence of intense strain fields. These distortions arise due to the lattice mismatch and dislocation accumulation at the interface, effectively lowering the energy barrier for martensite nucleation54. As such, these regions serve as ideal nucleation sites and provide a guiding framework for the formation and confined propagation of specific martensitic variants during cooling or mechanical loading. The combination of coherent crystallographic alignment and strain-concentrated interfacial zones enables a progressive, directional, and spatially uniform transformation process. This mechanism suppresses transformation-induced defects and enhances reversibility, ultimately contributing to the high fatigue resistance observed in the textured Ti49Ni51 alloy. The results highlight the critical role of microstructural design—particularly interface engineering—in achieving advanced functional properties in shape memory alloys.

Fig. 6: Lattice distortion near the B2/Ti₄Ni₂O interface revealed by HR-TEM.
Fig. 6: Lattice distortion near the B2/Ti₄Ni₂O interface revealed by HR-TEM.
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a Bright-field TEM image showing the boundary region between the B2 matrix and Ti₄Ni₂O precipitates. a′ SAED pattern taken across the B2/Ti₄Ni₂O interface along the [011]Ti₄Ni₂O // [011]B2 zone axis, confirming the epitaxial orientation relationship. b HR-TEM images focusing on the B2/Ti₄Ni₂O interface. c FFT reflections corresponding to b. d IFFT image generated by selecting the overlapping (100)B2 and (400)Ti₄Ni₂O reflections from the FFT in (b), clearly revealing the presence of interfacial dislocations. e–g Two-dimensional lattice strain fields near the interface are visualized using GPA based on the HR-TEM image in (b) including normal strain in the x-direction (ℇxx), y-direction (ℇyy), and shear strain (ℇxy), respectively, demonstrating significant local lattice distortion induced by the interface. The foot note “p” in a′ and c indicates signal from Ti₄Ni₂O precipitates.

Discussion

Achieving a simultaneous combination of a large elastocaloric effect and ultrahigh fatigue life has long been a central goal in the development of SMAs. In this study, we demonstrate that these typically conflicting properties can be reconciled through the introduction of epitaxially aligned Ti₄Ni₂O precipitates within a directionally solidified and textured B2 matrix. The resulting microstructure enables exceptional elastocaloric performance and mechanical durability, making the investigated alloy a promising candidate for next-generation solid-state refrigeration applications.

We propose a mechanism for the origin of the substantial elastocaloric effect observed in the alloy. First, the textured microstructure, fabricated via controlled directional solidification, enables a high average transformation strain of 6.1% under compressive loading—approaching the theoretical maximum of ~6.6% (see Fig. S3). This pronounced transformation strain is inherently linked to the crystallographic texture; prior studies have shown that SMA functionality is highly sensitive to grain orientation41,55. This is further evidenced by the significant anisotropy in superelastic behavior along and perpendicular to the SD, as presented in Fig. 2c. This anisotropic superelasticity and cyclic stability are consistent with both the stress–strain response and the microstructural evidence of preferential precipitate alignment (Fig. S9, S10). Compared with the perpendicular direction, compression along the SD benefits from the cooperative effect of <001 > B2 columnar texture and preferentially aligned Ti₄Ni₂O precipitates, which together lower the transformation stress, enhance transformation strain, and suppress defect accumulation. The enhanced superelastic strain along the SD indicates a higher degree of stress-induced martensitic transformability, thereby amplifying the elastocaloric response34. Second, the presence of epitaxially aligned Ti₄Ni₂O precipitates, sharing a fixed orientation relationship with the B2 matrix (Figs. 6 and S2), plays a critical role in suppressing dislocation activity. These precipitates impede slip-based plasticity and mitigate the accumulation of irreversible structural defects during cyclic loading56. As a result, the alloy is able to fully harness the potential of reversible SIMT to sustain large elastocaloric effects. Third, the columnar morphology of the Ti₄Ni₂O precipitates, aligned with the textured B2 grains along the SD (Fig. 1), supports an unobstructed SIMT within the textured microstructure. This alignment ensures minimal interference from secondary phases during transformation57,58. Together, these features allow the alloy to undergo nearly complete SIMT to B19′ martensite, yielding both high superelastic strain and a large adiabatic temperature change \({\Delta T}_{{ad}}\) during compression along the SD (Fig. 2c). In summary, the large elastocaloric effect originates from the alloy’s exceptional transformability, facilitated by its tailored microstructure.

The high fatigue resistance of the alloy is proposed to derive from two key aspects: mechanical stability (structural fatigue resistance) and transformation reversibility (functional fatigue resistance). As shown in Fig. 1a, b, the microstructure features linearly aligned Ti₄Ni₂O precipitates embedded within columnar B2 grains, both oriented along the SD. This mutual texture effectively acts as a “reinforced concrete” architecture during compressive loading, delaying crack initiation and propagation, and significantly enhancing fatigue resistance across millions of loading cycles. Compression loading reduces crack-driving forces and enhances structural fatigue resistance59; in the present alloy, this benefit is complemented by the reinforced architecture, together delivering ultrahigh operational life. More critically, in-situ characterization reveals that the transformation in the present alloy proceeds via numerous heterogeneous B19′ martensite nucleation events at B2/Ti₄Ni₂O interfaces, regulated by interfacial strain fields, which collectively result in a spatially macroscopic homogeneous transformation behavior in sharp contrast to the abrupt, avalanche-like B2–B19′ transformation with multi-twinned growth reported in conventional TiNi alloys50,60,61. The spatial confinement imposed by Ti₄Ni₂O precipitates limits transformation to nanoscale domains, suppressing martensite clustering and eliminating the characteristic stress plateau62,63. Consequently, this regulated pathway yields a nonlinear stress–strain response with narrow hysteresis, indicative of a progressive and highly reversible SIMT that remains stable over high fatigue cycles. The exceptional cyclic stability arises not from a second-order continuous transition64, but from a microstructurally regulated first-order pathway, where high-density Ti₄Ni₂O precipitates suppress strain localization and promote distributed nucleation, enabling a homogeneous, fatigue-resistant elastocaloric response. A slight initial \({\Delta T}_{{ad}}\) degradation (~2 × 10⁶ cycles), caused by dislocation accumulation and residual nanoscale martensite from lattice incompatibility (see Fig. S7), is constrained by the aligned precipitates and regulated progressive transformation. The resulting defects with nano martensite and dislocation substructures concentrated near precipitate boundaries are able to locally modify transformation stresses, leading to a sequential (two-step) macroscopic response that further enhances reversibility and fatigue resistance during extended cycling.

In conclusion, a textured TiNi alloy containing high-density Ti₄Ni₂O precipitates was successfully synthesized through directional solidification. The resulting columnar B2 grains are aligned along the solidification direction, with epitaxial Ti₄Ni₂O precipitates uniformly distributed throughout the matrix. A well-defined orientation relationship of {100} < 011 > B2 // {400} < 011>Ti₄Ni₂O is maintained across the matrix-precipitate interface. Under compressive loading along the SD, the alloy exhibits remarkable superelasticity with a recoverable strain of 6.2% and a low modulus of 38.3 GPa. Furthermore, it retains a large \({\Delta T}_{{ad}}\) of -15.9 K and a recoverable strain of 5.1% even after 10⁷ loading-unloading cycles. These outstanding properties are enabled by the synergy between a highly transformable textured B2 parent phase and the reversibility induced by well-aligned textured precipitates. The findings underscore the potential of textured TiNi alloys with epitaxial Ti₄Ni₂O precipitates as efficient and durable elastocaloric materials, and offer a pathway for future design of high-performance SMAs through the integration of microstructural texturing and precipitation engineering.

Methods

Sample presentation

Ti49Ni51 (at. %) ingots were synthesized via vacuum arc melting of high-purity Ti and Ni, with trace oxygen intentionally introduced to promote the formation of Ti₄Ni₂O precipitates. The oxygen contents of the starting Ti and Ni materials were measured by inductively coupled plasma mass spectrometry (ICP-MS) to be 0.061 wt.% and 0.007 wt.%, respectively, resulting in a total oxygen concentration of 0.029 wt.% in the as-fabricated alloy. To ensure chemical homogeneity, each ingot was re-melted five times prior to casting into a water-cooled copper crucible. Directional solidification was achieved by installing a water-cooled system beneath the crucible base, establishing a vertical thermal gradient during solidification. For comparison, a commercial polycrystalline Ti49Ni51 (at. %) alloy was also prepared. Specimens for both microstructural characterization and property evaluation were sectioned from the ingots using spark cutting.

Mechanical and elastocaloric testing

Compression tests were performed on cylindrical specimens with dimensions of 4 mm in diameter and 8 mm in height using an Instron 5969 universal testing machine. Tests were conducted at room temperature under a quasi-static loading-unloading strain rate of 10⁻³ s⁻¹, with strain tracked via a video extensometer. Fatigue performance was evaluated on an Instron 8801 fatigue test system under sinusoidal loading (10 Hz) for up to 10⁷ cycles, with a maximum stress of 900 MPa, consistent with the superelastic measurements. For elastocaloric measurements, \({\Delta T}_{{ad}}\) during unloading was captured by a K-type thermocouple welded to the specimen surface. Samples were preloaded to 900 MPa at a strain rate of 10⁻³ s⁻¹, held briefly for thermal equilibrium, and then rapidly unloaded at a strain rate of 0.2 s⁻¹ to ensure near-adiabatic conditions.

Structural and compositional analysis

All the samples for structural and compositional analysis were mechanically polished. SEM samples were etched using a mixture of 10% HF, 45% HNO₃, and 45% H₂O (by volume). TEM foils were prepared via twin-jet electropolishing in a solution of 20% H₂SO₄ and 80% CH₃OH (by volume) at 253 K. Phase identification was conducted using a Bruker D8 Advance X-ray diffractometer with Cu Kα radiation. Room-temperature microstructures were characterized using a FEI VERIOS 460 scanning electron microscope. Transmission electron microscopy investigations, including bright-field imaging, selected-area electron diffraction, energy-dispersive X-ray spectroscopy, and high-resolution TEM, were carried out on a JEOL JEM-F200 microscope. Crystallographic orientation and texture analyses were performed using a ZEISS Gemini 500 SEM equipped with an electron backscatter diffraction detector. EBSD data were processed using AztecCrystal software. Fast Fourier transformation, inverse fast Fourier transformation, and geometrical phase analysis were conducted using DigitalMicrograph software (version 3.43.3213.0).

In-situ transformation behavior analysis

Thermal transformation behavior was analyzed using a Netzsch 214 differential scanning calorimeter (DSC) across a temperature range of 153–263 K, at a constant heating and cooling rate of 10 K/min. To map strain evolution during compressive loading, digital image correlation was employed in conjunction with a high-speed camera. Tests were performed on rectangular specimens (3 × 3 × 8 mm³), loaded to 900 MPa at a strain rate of 10⁻³ s⁻¹ and unloaded to <10 MPa. In-situ loading XRD was conducted on a Rigaku MicroMax-007HF system equipped with a Mo rotating anode source (λ = 0.7093 Å). In-situ TEM cooling observations of martensitic transformation were carried out using a ThermoFisher Talos-F20 microscope, equipped with a Gatan 636 liquid nitrogen cooling holder. The TEM foil was held for 10 mins at each observation temperature to maintain thermal equilibrium.