Cryogenic superelasticity with large elastocaloric effect

Abstract

Shape-memory alloys suitable for low-temperature environments are crucially lacking despite their potential applications for aerospace engineering, superconducting technologies, and liquefied-gas-storage technologies. Even for benchmark shape-memory alloys such as Ti–Ni and Ni-based Heusler alloys, superelasticity becomes appreciably difficult to achieve upon cooling because of a drastic increase in stress hysteresis. Here, we report a superelastic strain of more than 7% for a Cu–Al–Mn shape-memory alloy down to 4.2 K with stress hysteresis as small as that at ambient temperature. Furthermore, the transformation entropy change remains as large as that at ambient temperature down to ~50 K. Consequently, an excellent elastocaloric cooling property is simultaneously obtained even at cryogenic temperatures, at which existing shape-memory alloys lose this property because of decreased entropy change and increased stress hysteresis. The developed cryogenic superelastic alloy shows excellent potential as a new class of material combining superelastic properties with sufficiently small dissipated energy and persistent elastocaloric cooling ability even in cryogenic environments.

Introduction

Shape-memory alloys (SMAs) are alloys that undergo the diffusionless solid–solid phase transformation known as martensitic transformation (MT). Their functionalities—the shape-memory effect and superelasticity—have long stimulated considerable interest in the field of materials science and engineering. The superelasticity enables a large recoverable strain to be attained using the stress-induced forward/reverse MTs via the application/release of uniaxial stress.¹ The behavior is similar to that of elastomers; the stress–strain (σ–ɛ) curve exhibits plateaus but not linearity as a superelastic response.¹ This functionality has resulted in many important applications, such as guide wires and dam** components.^{1, 2, 3} In addition, refrigeration using the latent heat yielded by the stress-induced MT has recently received considerable attention as an environmentally friendly refrigeration technique called elastocaloric cooling (EC),⁴ which may be a promising alternative to conventional vapor-compression techniques. This EC effect (previously termed the piezocaloric effect⁵) exploits the transformation entropy change ΔS as a cooling source under adiabatic conditions. ΔS is related to the equilibrium stress σ₀ and transformation strain ɛ_SE by the Clausius–Clapeyron equation:⁶

where V_m represents molar volume. Therefore, investigating the temperature dependence of superelastic behaviors can be used as an indirect approach to evaluate the EC potential of SMAs.

In the design of practical SMAs, minimizing the transformation hysteresis without deteriorating the shape-memory and/or superelastic property is a general strategy because it leads to reproducibility of these functionalities and enhancement of their fatigue properties.⁷ However, simultaneously achieving small σ_hys and large ɛ_SE (or ΔS) remains challenging. Even for the R-phase (trigonal-symmetric martensite) transformation in Ti–Ni, which exhibits a small temperature hysteresis of less than a few degrees Kelvin, the obtainable ɛ_SE is <1%;⁸ therefore, ΔS is also small. In contrast, σ_hys in SMAs with large ɛ_SE, for example, in FeNiCoAlTaB,⁹ is generally large. Recently, an extremely small temperature hysteresis and thereby excellent fatigue properties were achieved in certain SMAs^{10, 11, 12} because of the very small lattice mismatch between the parent and martensite phases^{10, 11} or among them and precipitates¹² arising from special relationships intrinsic to their lattice correspondence^{10, 11, 12, 13} (termed ‘supercompatibility’ by James¹⁴). Another, more conventional approach involves tailoring the microstructure to reduce intergranular constraints; the fatigue properties of Cu–Al–Mn SMA in the forms of a single crystal, bamboo-like polycrystals, and columnar-oriented polycrystals have been successfully improved while realizing small σ_hys and large ɛ_SE.^15,

Materials and methods

Ingots of Cu–17Al–15Mn and Cu–17Al–12Mn (at.%) were prepared by induction melting in an Ar atmosphere. The ingots were hot-rolled at 1143 K to reduce their thickness to ~3 mm. The obtained plates were subjected to cyclic heat treatment (see Supplementary Figure S1a) to obtain single crystals by inducing abnormal grain growth.²⁶ The plates (shown in Supplementary Figure S1b) were subsequently annealed at 473 K for 15 min to stabilize the MT behavior between the cubic (L2₁) parent and monoclinic (18R, now often termed 6M) martensite phases.^{27, 28} As shown in Supplementary Figure S2, it was confirmed that the obtained Cu–17Al–15Mn does not undergo the thermally induced MTs in the temperature range 4–300 K.

A series of tensile tests of Cu–17Al–15Mn single crystals was performed at a strain rate of 5 × 10⁻⁵ s⁻¹ in order from the highest to lowest test temperatures. To account for the training effect, the tests were conducted at least twice at every temperature, and the σ–ɛ curve obtained in the final test was used. The tensile orientations in the tested single crystals were determined using the electron backscattered diffraction technique.

Results and discussion

The tensile σ–ɛ curves of the Cu–Al–Mn single-crystal specimen (#1) in the low-strain region at temperatures of 4.2–160 K are presented in Figure 2. The series of σ–ɛ curves indicates that the superelastic property was successfully achieved at all the temperatures tested. The two insets show the full superelastic curve at 4.2 K and the crystallographic orientation of the tested crystal on an inverse pole figure. Those for another crystal (#2) with a different orientation are provided in Supplementary Figure S3. The ɛ_SE values of 7.1% and 5.3% for specimens #1 and #2 at 4.2 K are less than the theoretical values of 9.1% and 7.6%, respectively, estimated from the crystallographic relationship between the cubic (L2₁) parent and monoclinic (18R) martensite phases in a Cu–17Al–10Mn SMA at ambient temperature.²⁹ This discrepancy may have been caused by a change in the lattice parameters resulting from the different Mn concentration and temperature. The critical stresses, σ_Ms and σ_Af, of the stress-induced forward/reverse MTs, as depicted in Figure 2, are plotted in Figure 3a as a function of temperature. Those reported for bamboo-textured³⁰ and columnar-grained²⁷ high machinability,²⁷ improved fatigue life,¹⁶ grain size tunability,²⁶ and transformation temperature tunability.²⁸ It has been reported that the near <100>-oriented single crystal exhibits very large ɛ_SE approaching 10%.³¹ If we assume ɛ_SE=10%, σ_Ms=50 MPa at 4.2 K, temperature-independent σ_hys=20 MPa, and critical slip stress=500 MPa in the near <100>-oriented single crystal, operations of superelasticity at temperatures up to 366 K and of the EC effect at temperatures as low as T_CO=24 K may be feasible; one specimen can exhibit both superelasticity and the EC effect over a wide temperature window of ~340 K. As previously mentioned, the crystallographic compatibility between the parent and martensite phases^{10, 11, 12, 13} was considered to be a factor determining the magnitude of the MT temperature hysteresis. Although experimental investigations have not been attempted, this guidance appears to be valid for suppressing σ_hys and temperature hysteresis; thus, the existing COP_mat metrics (Figure 5) may be revised for such supercompatible superelastic alloys. However, it is concerning that these conditions can be easily disrupted by changing temperature and composition. It remains unknown how thermal activation develops as temperature decreases because the extremely small hystereses are all observed at ambient temperatures at which the athermal nature is dominant. Therefore, the novel feature of the present Cu–Al–Mn alloy, namely, the temperature-independent small σ_hys and persistent EC performance over a very wide temperature range, will not be lost even after the supercompatible strategy is applied to design a superelastic alloy.

As shown in Figure 6, the present Cu–Al–Mn alloy exclusively covers the application fields of space engineering, superconducting technologies, and liquefied-gas technologies, for which it has been almost impossible to implement existing SMAs. This material will contribute to innovative developments of these cryogenic technologies as well as the downsizing or simplification of various cryogenic systems.

Figure 6

Existing and potential application fields and operational temperature windows of shape-memory materials.