Quasi-superplasticity in the AlCoNiV medium entropy alloy with Heusler L2₁ precipitates Open Access

The tendency of materials to undergo extreme plastic deformations (i.e., >400% failure elongation) at high temperatures without notable necking or fracture is widely termed superplasticity.¹ As a prerequisite, fine grain sizes, typically less than 10 μm, and relatively high testing temperatures close to or near ∼0.5 T_m, where T_m is the absolute melting temperature of the material, are required for superplasticity.² Exposure to high deformation temperatures where grain growth is inevitable is a major challenge since the rapid grain growth inhibits superplastic flow.^1,2 As a result, superplasticity is ideal in dual and multi-phase alloys where the secondary phases can inhibit grain growth. The underlying mechanisms for achieving the high superplastic elongations are either dynamic recrystallization (DRX), which is a recrystallization phenomenon that occurs under the action of deformation, or grain boundary sliding (GBS).³ The principle of superplasticity is used by the metal-forming industries to form complex geometric parts from sheet metals for applications in the automotive and aerospace industries.³ Ideally, the term superplasticity applies to alloys with high strain-rate sensitivity (m) > 0.5; however, there are a few coarse-grained materials, especially Al–Mg-based alloys, which achieve elongations of up to 200%–300% at high temperatures.⁴ The strain rate sensitivity (m) of these alloys is about 0.33, and their higher elongations are termed “extended ductility” or “quasi-superplasticity.”^4,5 These levels of elongation are also acceptable for industrial forming processes, such as extrusion and stamping.

Medium and high entropy alloys (M/HEAs) have garnered widespread attention in recent years.^6,7 Although composed of multiprincipal elements, HEAs and medium entropy alloys (MEAs) typically form simple solid solutions of either face-centered cubic (fcc), body-centered cubic (bcc), or hexagonal closed-packed (hcp) structures instead of complex intermetallics by utilizing their inherently high configurational entropies to overcome the enthalpies of compound formation.^6,7 M/HEAs exhibit an exciting range of physical and functional properties, such as high strength and fatigue resistance,^8,9 excellent cryogenic damage tolerance,^10,11 corrosion resistance,^12,13 resistance to hydrogen embrittlement,^14,15 high-temperature strength,^16,17 soft magnetic properties,^18,19 and others.^20,21 Based on the requirements for superplastic flow, the ability of M/HEAs to display superplastic forming abilities would present an important breakthrough in materials science. This would mean that these M/HEAs could be processed at an industrial scale so they could finally fulfill their commercialization potential.

Kuznetsov et al.²² reported the first evidence of superplastic flow in the multiphase AlCuCrFeNiCo HEA. The HEA displayed maximum elongations of up to 860% at 1273 K and was enabled by GBS in the fine-grained fcc and bcc structures. In another study, a maximum elongation of 1240% was observed at 1273 K in the AlCuCrFeNiCo HEA.²³ Nguyen et al.²⁴ observed a superplastic flow (600%–770%) between 973 and 1073 K in the V₁₀Cr₁₅Mn₅Fe₃₅Co₁₀Ni₂₅ HEA processed by high-pressure torsion (HPT) with an average grain size of ∼30 nm. The superplasticity response was attributed to grain boundary sliding, which maintained the equiaxed structure in both deformed and undeformed regions of the tensile specimen. HPT was also used to produce ultrafine grains in the Cantor alloy, which resulted in >600% superplastic flow at 973 K, which was attributed to grain boundary sliding (GBS).²⁵ By introducing B2 precipitates into the Cantor alloy via Al-alloying and HPT processing to obtain ultrafine grains, a high-strain rate superplasticity of ∼2000% was observed at 1073 K due to GBS and the strong pinning effect of the B2 grains, which restricted grain growth.²⁶ Very recently, the fcc-B2-σ triplex Al_0.3CoNiCr MEA was also reported to undergo a high-strain rate superplasticity of ∼1175% at 1073 K via GBS and partial melting of pure Al at grain boundaries/interfaces.²⁷ 

From the above-mentioned evidence of superplasticity in M/HEAs, it is seen that severe plastic deformation (SPD) techniques, such as HPT and forging, are the most common techniques used for highly refined grain structures suitable for superplasticity. However, investigations about superplasticity in conventional thermomechanical processed M/HEAs are very limited. The earlier work by Reddy et al.,²⁸ where the annealed fcc-structured CrMnFeCoNi HEA with 1.4 μm grain size displayed ∼320% ductility at 1023 K via GBS, and the recent work by Sohn et al.,²⁹ where the annealed CoNiV with κ/σ micro-duplex structure displayed ∼450% high-strain rate superplasticity at 1073 K due to enhanced DRX effects, are the most notable examples of superplasticity in conventionally processed M/HEAs. Inspired by this, we develop the Al₇(CoNiV)₉₃ (at. %) MEA with a duplex microstructure of fcc-matrix and bimodal L2₁ Heusler precipitates via conventional cold-rolling + annealing. The alloy displays a high fracture elongation of ∼270% at 1073 K, which is enabled by DRX effects that induce the formation of an equiaxed fcc-L2₁-σ triplex microstructure. The estimation of the strain-rate sensitivity (m) of the studied alloy is found to be ∼0.32 in agreement with conventional alloys, which undergo quasi-superplasticity at elevated temperatures.

Al₇(CoNiV)₉₃ (at. %) alloy ingots were prepared by arc-melting high-purity Al, Co, Ni, and V elements (>99.9 at. %). Each ingot was remelted at least five times to ensure chemical homogeneity, after which they were suction-cast into an 80 × 10 × 2.5 mm³ copper mold. The sheet samples were then sealed in vacuum-pumped quartz tubes, homogenized at 1373 K for 8 h, and then quenched in water. The homogenized sheets were cold-rolled to ∼65% thickness reduction, then subsequently annealed at 1173 K for 1 h, and quenched in water.

The phases of the samples were determined by the Rigaku SmartLab^® x-ray powder diffractometer (XRD) using a Cu target and a step scan of 2°/min operated at 20 kV. The microstructural grain size and morphology were investigated using electron back-scattering diffraction (EBSD) and scanning electron microscopy (SEM). Samples for the EBSD studies were mechanically polished using several grit sizes and a final polishing stage of colloidal silica for 1 h. EBSD scans were also performed using a step size of 60 nm. SEM/EBSD characterizations were carried out using a field emission Zeiss Supra 55 Scanning Electron Microscope. Thin transmission electron microscopy (TEM) foils were prepared using the FEI Strata 400S Dual-Beam focused ion beam (FIB). The Tecnai G2 F20 S-TWIN FEI TEM operating at an acceleration voltage of 300 kV was used for microstructural investigations of the thin foils.

Flat dog-bone-shaped tensile specimens of dimensions 8.5 × 2.2 × 0.85 mm³ were cut from the annealed samples by electrical-discharge machining. The CMT5205 SANS universal testing machine was used to perform uniaxial tensile tests from room to elevated temperatures at different strain rates with an extensometer attached to estimate the strain. Before every test, the stability of the testing temperature measured by a thermocouple was ensured by heating the sample to the corresponding temperature and holding it for about 30 min. The reproducibility of the tests was ensured by repeating each test three times. Strain-rate jump tests were also performed to estimate the strain-rate sensitivity by using a varied strain rate between 10⁻³ and 10⁻¹ s⁻¹.

Figure 1 shows the microstructure of the prepared Al₇(CoNiV)₉₃ MEA. The XRD pattern [Fig. 1(c)] shows that the alloy consists of dominant fcc peaks with minor bcc peaks, indicating a dual-phase microstructure. The {111} and {200} superlattice peaks at ∼26.6° and ∼30.9° that are known to belong to the ordered (Co, Ni)₂AlV-L2₁ bcc phase are also found.³⁰ The lattice parameters of the fcc and L2₁ phases are estimated from the XRD peaks to be 3.612 and 5.811 Å, respectively. The EBSD phase and inverse pole figure (IPF) maps in Figs. 1(a) and 1(b) show a fully recrystallized duplex microstructure with an fcc matrix dispersed with L2₁ precipitates. The volume fractions of the matrix and precipitate phases are about 80% and 20%, respectively. The fcc matrix has an average grain size of ∼1.4 μm, while the L2₁ precipitates show a bimodal distribution ranging from the nanometer to the micrometer scale. The L2₁ grains exist as coarse spheroid (∼8 μm in diameter) and rod-like shaped particles ranging from several hundreds of nanometers to tens of micrometers with fine nanometer-sized fcc laths emerging from within these regions. The rod-like shaped L2₁ precipitates have an average aspect ratio of ∼2.5. In addition, the ultrafine L2₁ grains are also distributed mostly at the triple junctions of fcc grains. The bimodal distribution of the L2₁ grains is attributed to its rapid agglomeration tendency³¹ and originates from the high energy shear bands induced during cold-rolling, which act as heterogeneous nucleation sites during recrystallization annealing.³⁰ The unique bimodal distribution of the semi-coherent L2₁ precipitates in the fcc matrix allows for efficient precipitation strengthening.³⁰ The fcc/L2₁ duplex microstructure is also shown in the BF-TEM image in Fig. 1(d). The fcc crystal structure of the matrix grains and the ordered bcc-L2₁ crystal structure of the precipitates are confirmed by their corresponding selected area diffraction (SAED) images in Figs. 1(e) and 1(f). In terms of chemical composition, the sample reveals a clear discrepancy in the Al content, namely the Al-deficient matrix and the Al-rich precipitate regions (Fig. S1).

Figure 2(a) shows the representative tensile engineering stress–strain curves at different temperatures of the studied alloy at a constant strain rate of 10⁻³ s⁻¹. At room temperature, the alloy shows a yield strength of ∼1190 MPa, an ultimate tensile strength of 1585 MPa, and fracture elongation of ∼32%, which is similar to the fully recrystallized fcc/L2₁ duplex Al_0.2CoNiV MEA.³⁰ As the temperature is increased to 923 K, the alloy exhibits a yield strength of 1030 MPa, a tensile strength of 1090 MPa, and fracture elongation of ∼17%. Here, the alloy still undergoes a significant amount of strain-hardening (1090–1030 MPa = 60 MPa), although very small compared to the room temperature value of 395 MPa. Due to the chemical complex nature of M/HEAs, the precipitate–matrix interactions can generate high lattice distortions to increase dislocation slip resistance beneficial for high strength even at elevated temperatures.³² The ability of the alloy to preserve excellent high-temperature strength and strain hardening capability is expected to be utilized for high-temperature applications. Furthermore, a transition into flow-softening tensile behavior is seen to occur at 973 K, with a fracture elongation of ∼70% and a peak flow stress of 875 MPa. As the temperature increases from 1023 to 1073 K, the peak flow stress decreases from 475 to 234 MPa, while the fracture elongation increases from ∼150% to ∼270%. The strain rate dependence of the alloy at 1073 K is shown in Fig. 2(b). As the strain rate increases from 10⁻³ to 10⁻¹ s⁻¹, the peak flow stress increases from 234 to 820 MPa, while the elongation decreases from 270% to ∼105%. The increase in flow stress with increasing strain rates at high temperatures is typical of alloys that undergo superplastic deformation.^24,29

To understand the high-temperature deformation response of the alloy, we estimated the strain-rate sensitivity (m) and activation energy (Q) of the alloy. These parameters are determined from the Zener–Hollomon equation,³³ $Zσ=Aσn=ε̇exp(QRT)$⁠, where $ε̇$⁠, Q, R, T, n, σ, and A are the strain rate, activation energy, molar gas constant, absolute deformation temperature, stress exponent, plastic flow stress, and a deformation mechanism sensitive constant, respectively. The strain-rate sensitivity (m) is also defined as m = $∂⁡ln⁡σ/∂⁡lnε̇$ and can be estimated from the linear ln σ − ln $ε̇$ plot in Fig. 2(c). Here, the strain-rate sensitivity is calculated to be 0.32 with a correlation coefficient (R²) of 97.6%, suggesting a satisfactory fit for the Al₇(CoNiV)₉₃ MEA. The activation energy $Q=nR∂⁡lnσ∂(1/T)$ provides information about the ease at which diffusion or superplasticity proceeds in an alloy and can be determined from Fig. 2(d). Based on the $∂⁡lnσ∂(1/T)$ slope and the satisfactory correlation coefficient (R²) of 97.9%, the stress exponent n, which is the inverse of the strain rate sensitivity, is estimated as n = 1/m = 3.215. The activation energy for Al₇(CoNiV)₉₃ is then determined to be ∼421 kJ/mol. This shows that the activation energy of Al₇(CoNiV)₉₃ is about 2.5 times higher than the activation energy (161–182 kJ/mol) of the CoNiV alloy and also higher than the grain boundary diffusion values of the individual elements.²⁹ The dependence of strain rate sensitivity on strain rate and the tensile strain was estimated by conducting strain-rate jump tests at 1073 K shown in Fig. S2. The m value shows a decrease from 0.89 at 0.001 s⁻¹ to 0.41 at the largest strain rate of 0.01 s⁻¹ and then decreases to 0.34 at 0.001 s⁻¹. The variation in m value suggests a strong microstructural sensitivity at different strain rates and ultimately determines whether or not superplasticity will occur.³⁴ 

To understand the influence of temperature on the present phases in the studied alloys, we performed room-temperature XRD measurements on the grip portion (i.e., undeformed region) of the tensile samples. Figure 3(a) shows the XRD patterns of Al₇(CoNiV)₉₃ deformed at 973, 1023, and 1073 K at a strain rate of 10⁻³ s⁻¹. At 973 K, the alloy is still composed of an fcc/L2₁ duplex structure. By increasing the testing temperature, a small fraction of the σ phase appears in the samples tested at 1023 and 1073 K. The strain-rate dependence of the phase evolution at 1073 K is shown in Fig. 3(b). The time taken for each tensile test at different strain rates is depicted in this figure. After only 22 s of tensile testing, the undeformed region still shows predominant duplex fcc/L2₁ peaks. As the deformation time increases from 120 to 2862 s (i.e., strain rate decreases), the sigma peaks appear, resulting in the development of an fcc-L2₁-σ triplex in the undeformed region. The increase in the sigma amount at lower strain rates is because the sigma phase formation is a diffusion-controlled phase decomposition process, which is more feasible at longer periods.³⁵ This is a characteristic of high V containing M/HEAs.³⁶ The formation of the brittle σ phase at these temperatures typically aids superplastic flow by inhibiting grain growth of the parent phase(s), which is suitable for GBS.^24,37

The effects of strain rate on the superplastic deformation response of the alloy at 1073 K were also probed in the tensile gage by EBSD represented in Fig. 4. The microstructures that have developed here post-deformation are completely different from the initial microstructure. At faster strain rates, significant refinement for both fcc and bcc-type grains is detected, which have developed into lamellar-like structures that are elongated in the tensile direction [Figs. 4(a) and 4(b)]. Several cracks are also formed in the high-strain rate samples; hence, the corresponding fracture elongations are also reduced. By decreasing the strain rates, the lamellar-like grain development diminishes, and the microstructure transitions into a near-equiaxed type that usually favors superplastic flow, as seen by the corresponding higher fracture elongations at lower strain rates in Fig. 2(b).

Figure 5 shows the microstructural evolution of the tensile specimen deformed at 1073 K at 10⁻³ s⁻¹ strain rate, which showed the largest elongation of ∼270%. Figure 5(a) shows the EBSD IQ image overlayed with the phase maps at the grip part of the tensile sample. Here, the microstructure is akin to that of the initial microstructure with fcc grains and bimodal L2₁ grains. However, in addition, σ grains are also formed, indicating a transition into a triplex fcc-L2₁-σ microstructure. The fcc grains also undergo slight grain growth, developing an average grain size of 2.7 µm. This slight increase in grain size at the tensile grip also signifies time-dependent static grain growth at this temperature. The volume fractions of the fcc, L2₁, and σ phases are found to be 66%, 22%, and 12%, respectively. At the grip/gage interface [Fig. 5(b)], the volume fractions of the σ and L2₁ phases increase at the expense of the fcc phase. Figure 5(c) shows significant grain refinement in the initially existing fcc and L2₁ grains at the uniform gage region. Finally, the near fracture region of the tensile specimen shows an equiaxed multiphase microstructure developed under the action of deformation. The volume fractions of the comprising phases in this region are 8% fcc, 41% L2₁, and 51% σ phases. The average grain sizes of the L2₁ and σ grains are about 800 and ∼700 nm, respectively, while the fcc grains are mostly below 500 nm. The variation in microstructures at the tensile gage compared to the grip and initial microstructures indicates the occurrence of dynamic recrystallization (DRX).^28,29 The disappearance of the initial coarse L2₁ grains into refined spherical equiaxed grains is very crucial to the transition into superplastic-like deformation. Aided by the refinement of the microstructure and the formation of the σ grains, the interfacial boundaries can slide to accommodate stress during the deformation.

To confirm the existing phases at this temperature during static and dynamic deformations, TEM investigations were conducted. Figure 6(a) shows the STEM and corresponding energy dispersive x-ray spectroscopy (EDS) elemental maps of the sample taken from the grip region of the tensile sample deformed at 10⁻³ s⁻¹ and 1073 K. Here, the microstructure is akin to the initial undeformed sample with dominant fcc grains with secondary precipitates. The EDS elemental maps show the Al-rich regions as the notable compositional variation between the matrix and precipitates. Figure 6(b) shows the STEM image of the near fracture region of the tensile specimen. The alloy shows a microstructure of equiaxed grains. The corresponding EDS elemental maps show profound Al and V partitioning in certain grains. The Al-rich grains are confirmed by SAED imaging to correspond to the L2₁ phase while the V-rich grains belong to the σ phase. The recrystallized triplex microstructure at the near fracture region shown here agrees with the EBSD results. The formation of the V-rich σ phase under static and dynamic evolutions at 1073 K is expected since the fcc phase undergoes phase decomposition to form either the κ or σ phases below 1173 K in the CoNiV alloy system.³⁵ The κ and σ phases are easily distinguished by the high V amount in the σ phase;²⁹ hence, the κ phase was not found in the studied alloy system.

High-temperature deformation where flow softening occurs to favor higher elongations can be accommodated by dislocation glide, resulting in the elongation of grains along the stress direction or DRX-induced GBS to accommodate dislocation climb.^2,37 The latter usually results in the simultaneous refinement and rotation of grains to achieve superplastic flow, especially when grain growth is immensely suppressed. In the currently studied alloy, the dynamic evolution shows DRX occurs at 1073 K at a strain rate of 10⁻³ s⁻¹. While the L2₁ Heuser phase is known to be subject to rapid grain growth,³¹ the reversion from the bimodally distributed nanoscale L2₁ grains and coarse L2₁ islands to a rather equiaxed grain distribution implies that DRX has occurred. Furthermore, we believe that under the action of stress, the refinement of the microstructure is mainly assisted by the formation of the harder σ phase, which restricted the grain growth of the Heusler grains. Regardless of having a highly refined microstructure, the alloy displays only 270% elongation, which is characterized as quasi-superplasticity.¹ 

In conventional alloys, a strain rate sensitivity of ∼0.33 is related to the solute drag creep behavior where solute atoms segregate favorably at dislocations and are dragged by the moving dislocations to achieve elongations ∼200%–300%.^1,38 While this behavior ideally occurs at an m of ∼0.33, the dislocations can separate from the atmosphere of the solute atoms when they transition from glide to climb, thus increasing the strain rate sensitivity to ∼0.5 to favor superplasticity.^5,39 However, in multicomponent alloys, distinguishing between solute and solvent atoms sometimes is a challenge. Owing to the chemical composition of the studied alloy, the Al atoms due to their limited amount and large size can be considered solute atoms. In addition, since we estimated the strain-rate sensitivity and stress exponent of the alloy to be 0.32 and 3.125, the likelihood of viscous glide occurring at the early stages of the deformation cannot be ruled out due to the tendency of the solute Al atoms to segregate at the misfit phase boundaries where dislocations exist.⁴⁰ However, due to the diminished stability of the fcc phase at this temperature at the expense of the σ phase, refinement of the microstructure occurs as deformation progress. This results in the development of the equiaxed triplex fcc-L2₁-σ microstructure. Although the observed DRX effect should be favorable for larger superplastic elongations, the ∼270% fracture elongation displayed by Al₇(CoNiV)₉₃ in this work is noticeably smaller than that of CoNiV (∼415%) at the same strain rate.²⁹ This is because the σ amount formed here is very high (∼50%) and since they are unable to accommodate dislocations during GBS, they increase cavitation, which leads to the early fracture of the sample.^26,41 In addition, ductility at high temperatures can also be inhibited by the surface oxidation of the deformation samples. Considering the oxidation problems presented by vanadium, the formation of the V-rich sigma phase at this temperature makes the alloy susceptible to immense oxidation, as evidenced by the green and black coloration of the sample (Fig. 5). Therefore, the observed quasi-superplastic behavior at this temperature is mostly dominated by the DRX effect, which causes the formation of the refined triplex fcc-L2₁-σ microstructure evidenced by the flow softening. Even though this quasi-superplastic elongation of ∼270% is still adequate for industrial superplastic forming operations,⁴² grain refinement techniques could be used to shift the alloy into higher superplastic elongations. Therefore, the present study highlights the need for future investigations into the cavitation and oxidation characteristics to understand the creep properties and the overall high-temperature deformation mechanisms of this alloy and other Heusler-type M/HEAs.

In summary, the microstructure and tensile deformation behavior of the Al₇(CoNiV)₉₃ MEA are investigated with emphasis on the high-temperature tensile behavior and the underlying microstructural evolution. Thermomechanical treatment of the Al₇(CoNiV)₉₃ MEA resulted in the formation of a dual-phase fcc and an Al-rich L2₁ Heusler-ordered bcc hierarchical microstructure. The alloy consists of fine fcc grains as the matrix and L2₁ grains as grain boundaries, triple junctions, and coarse micron-scale islands. The high-temperature tensile tests reveal >1 GPa yield strength at 923 K with good strain hardening capacity, resulting in a good ductility of ∼17%. Transition into flow softening deformation responses occurs at temperatures above 973 K, with a maximum fracture elongation of ∼270% observed at 1073 K. Strain rate tests showed the flow strength increased at high strain rates while the ductility decreased with increased strain rate. The strain-rate sensitivity and stress exponent of the alloy were estimated to be 0.32 and 3.125, which are typical of alloys that show quasi-superplastic behavior or extended ductility (200%–300%). Microstructural investigations revealed dynamic recrystallization (DRX) and formation of the V-rich sigma phase to be responsible for the quasi-superplasticity displayed by the alloy. The inability of the alloy to meet the criterion for superplasticity was attributed to the high volume fraction of the sigma phase, which increased cavitation. However, this quasi-superplasticity elongation is still acceptable for industrial superplastic forming operations. This shows that by conventional thermomechanical treatment strategies, alloys could still be processed by metal forming operations, although further grain refinement strategies can be used to achieve higher superplastic elongations.

This section contains two additional figures captioned as Figs. S1 and S2Figs. S1 and S2, respectively.

This work was supported by the National Natural Science Foundation of China (Grant Nos. U1832203, 11975202, and 12275237), the Natural Science Foundation of Zhejiang Province (Grant Nos. LY15E010003 and LZ20E010002), and the Fundamental Research Funds for the Central Universities.

The authors have no conflicts to disclose.

Raymond Kwesi Nutor: Data curation (equal); Formal analysis (equal); Writing—original draft (equal); Writing—review & editing (equal). Ran Wei: Data curation (equal); Investigation (equal); Methodology (equal). Qingping Cao: Investigation (equal); Methodology (equal); Supervision (equal); Writing—review & editing (equal). Xiaodong Wang: Formal analysis (equal); Methodology (equal); Validation (equal). Shaoqing Ding: Methodology (equal); Resources (equal). Dongxian Zhang: Project administration (equal); Resources (equal); Software (equal). Fushan Li: Methodology (equal); Validation (equal); Writing—review & editing (equal). Jian-Zhong Jiang: Conceptualization (equal); Funding acquisition (equal); Supervision (equal); Writing—review & editing (equal).

The data that support the findings of this study are available from the corresponding author upon reasonable request.