Abnormal Effect of Al on the Phase Stability and Deformation Mechanism of Ti-Zr-Hf-Al Medium-Entropy Alloys

Abstract

Complex concentrated alloys, including high-entropy alloys (HEAs) and medium-entropy alloys (MEAs), offer another pathway for developing metals with excellent mechanical properties. However, HEAs/MEAs of different structures often suffer from various drawbacks. So, investigations on the effect of phase and microstructure on their properties become necessary. In the present work, we adjust the phase constitution and microstructure by Al addition in a series of (Ti₂ZrHf)_100−xAl_x (x = 12, 14, 16, 18, 20, at.%, named Al_x) MEAs. Different from traditional titanium, Al shows a β-stabilizing effect, and the phase follows the evolution of α′(α)→α″→β + ω + B2 with Al increasing from 12 to 20 at.%, which could not be predicted by the CALPHAD (Calculate Phase Diagrams) method or the Bo-Md diagram because of the complex interactions among composition elements. At a low Al content, the solid solution strengthening of the HCP phase contributes to the extremely high strength with a σ_0.2 of 1528 MPa and σ_b of 1937 MPa for Al14. The appearance of α″ deteriorates the deformation capability with increasing Al content in the Al16 and Al18 MEAs. In the Al20 MEA, Al improves the formations of ordered B2 and metastable β. The phase transformation strengthening, including B2 to BCC and BCC to α″, together with the precipitation strengthening of ω, brings about a high work-hardening ratio (above 5 GPa) and improvements in ductility (6.8% elongation). This work provides guidelines for optimizing the properties of MEAs.

1. Introduction

To date, high-entropy alloys (HEAs) or complex concentrated alloys (CCAs) based on multi-principal elements have fascinated many researchers [1,2,3] because of their infinite compositional space and excellent properties such as good mechanical properties [4], high corrosion resistance [5], and radiation resistance under extreme conditions. The mechanical response of this kind of alloy is remarkably dependent on the phase and microstructure. Most research studies focus on FCC or BCC HEAs, which have unique advantages but also show some inherent drawbacks. Face-centered cubic (FCC)-structured HEAs based on the traditional Cantor alloy (FeCoNiCrMn) show good deformation ability but low yield strength at room temperature, while body-centered cubic (BCC) HEAs, including refractory elements, exhibit high strength and insufficient ductility. As one of the important common crystal structures, the hexagonally close-packed (HCP) structure is common in Mg-, Al-, and Ti-based alloys. However, HCP-structured HEAs suffer from intrinsic brittleness without any tensile ductility, have low strength less than 500 MPa, such as Al_19.9Li₃₀Mg₃₅Si₁₀Ca₅Y_0.1 [6] and Hf₂₅Sc₂₅Ti₂₅Zr₂₅ [7], or contain lots of rare-earth elements, such as GdTbDyTmLu [8] and YGdTbDyLu [8].

Recently, in pursuit of superior performance, research studies on HEAs have not only focused on stable or single-phase systems, but HEAs with complex phase compositions, including dual-phase like FCC + L1₂ [9], BCC + B2 [10], and metastable HEAs like Ti₃₅Zr_27.5Hf_27.5Nb₅Ta₅ [11], TiZrHfTa_x [12], Fe_80−xMn_xNiCr [13], and eutectic HEAs like Ni_2.1FeCoCrAl [14] have received much attention because of their unique properties and deformation mechanism. The adjustment of phase constitution and stability is the key point for developing new kinds of HEAs with special objects. One of the most efficient methods to change the microstructure is to adjust its composition.

As an important alloying element in HEAs, the effect of Al on the phase transition and mechanical properties has been explored in many different systems. For example, Al brings in phase evolvement from FCC or HCP to a BCC crystal structure in HEAs such as Al_xCoCrFeNi(Mn) [15,16], Al_xCrMnFeCoNi [17], and CoFeMnNiMo_0.2Al_x [18]. It is proposed that the dissolution of Al in FeCoNiCrMn would distort the lattice and increase the lattice constant, resulting in a large lattice distortion energy that promotes the transformation from FCC to BCC, which is originally formed by the disordered BCC (A2) phase and the ordered BCC (B2) phase. Al_xCrMnFeCoNi goes through the process of FCC to BCC to the nano-precipitated A2 phase. In CoFeMnNiMo_0.2Al_x, the HCP structure gradually changes to the L1₂ phase and the BCC phase because of the mixed action of Al and Mo, where Al was proven to be an effective stabilizer of the BCC phase. Al also improves the formations of B2 and the Laves phase in refractory HEAs like TiZrHfNbAl_x [19], TiZrHfNbTaAl_x [20], TiZrHf_0.4NbTaAl_0.4 [10], TiZrNbTa_0.5MoAl [10], Ti-Ta-Zr-Al_x [21], and NbTa_0.5TiAl_x [22], which mostly consist of a single BCC phase. In the Ti-Ta-Zr-Al system, the transition from BCC to the ordered B2 phase occurs when the Al content is greater than 5%. It is worth mentioning that the increase in Al decreases the lattice constant in TiZrHfNbTaAl_x [20]. In NbTa_0.5TiAl_x, it was thought that a small amount of Al could promote the formation of the BCC phase and an excess of Al would lead to the generation of the B2 phase. In traditional titanium alloys, Al acts as an α-stabilizing and solid solution element for strengthening and improving oxidation resistance. However, it is also reported to be able to improve the formation of the β phase in the Ti-Zr system [23]. So, it is possible to modify the phase and microstructure in Ti-containing HEAs and achieve outstanding mechanical properties by alloying with Al.

In the present work, we adjust the phase constitution and microstructure of (Ti₂ZrHf)_100−xAl_x MEAs by Al addition. The phase evolution and mechanical properties of these MEAs are investigated. Based on this, the effect of Al on the phase composition is discussed and compared with the traditional phase prediction method. The deformation mechanism determined by the phase constitution and stability is also discovered. This work provides guidelines for the optimization of the composition and properties of MEAs, especially ultra-strength MEAs.

2. Materials and Methods

In this work, (Ti₂ZrHf)_100−xAl_x MEAs (denoted as Al_x, x = 12, 14, 16, 18, and 20, at.%), compositions shown in

Table 1. The compositions of the (Ti₂ZrHf)_100−xAl_x MEAs (Al_x, x = 12, 14, 16, 18, and 20, at.%).

Alloys	Al (at.%)	Hf (at.%)	Ti (at.%)	Zr (at.%)
Al12	12	22	44	22
Al14	14	21.5	43	21.5
Al16	16	21	42	21
Al18	18	20.5	41	20.5
Al20	20	20	40	20

, were prepared by arc melting the elements with high purity inside a chamber filled with low-pressure high-purity argon atmosphere. To ensure chemical homogeneity, the ingots were flipped over and re-melted at least 5 times. After that, the melt was drop-casted into a water-cooled copper mold with a size of 10 mm × 10 mm × 50 mm.

The phase constitution of the specimens was characterized by X-ray diffraction (XRD) with Cu K_α radiation in a 2θ range from 20° to 100° at a rate of 6°/min. The evolution of the microstructure with Al was investigated by using a Zeiss Supra55 scanning electron microscope (SEM, Carl Zeiss AG, Oberkochen, Germany). The SEM samples were ground with SiC abrasive papers followed by mechanical polishing. Then, the samples were etched with Kroll’s reagent (HF:HNO₃:H₂O = 1:3:7 in volume). A more detailed microstructure characterization was conducted by a transmission electron microscope Tecnai F30 (TEM, FEI Company, Hillsboro, America). The TEM samples were cut from a Φ3 mm cylinder with a thickness of 0.3 mm, followed by mechanical grinding to 50–60 µm, and then jet polishing at −35 °C using double spray liquid with methanol–perchloric acid–glacial acetic acid = 88:8:4 in volume.

Tensile tests were carried out at room temperature using a CMT4105 universal electronic tensile testing machine (SNSCK Company, Shenzhen, China) at a strain rate of 1 × 10⁻³ s⁻¹ on a plate specimen with a gauge length of 12 mm, a width of 3 mm, and a thickness of 2 mm. The strengthening features were evaluated by work hardening rate curves, which were calculated by differentiating the true stress–strain curves (dσ/dε). The compression test was conducted on the cylindrical sample with a diameter of 6 mm and a height of 9 mm at the loading rate of 5 × 10⁻⁴ s⁻¹.

The phase compositions of Al_x MEAs were calculated by using Thermo-Calc 2022 with the TCTI3 database at the temperature range from 20° to 100°.

3. Results

3.1. Phase Prediction by CALPHAD and Traditional Bo-Md Method

The phase evolutions of Al_x MEAs predicted by CALPHAD are shown in Figure 1. It can be seen that the melting points decrease from 1478 °C to 1325 °C with the increase in Al content. For Al12 to Al16, the body-centered cubic BCC_B2 phase firstly solidifies from the liquid phase and then transfers to the hexagonal close-packed HCP_A3 phase together with the remaining liquid phase. For Al18 and Al20, the HCP_A3 phase directly solidifies from the liquid phase. Except for Al20, there are single HCP_A3 regions in the other four alloys. With decreasing temperature, different kinds of Zr-Al intermetallic form, and changes occur from Al₃Zr₅ to AlZr₂ for Al12 and Al14, Al₃Zr₅ to Al₂Zr₃ for Al18 and Al20, and AlZr₂ to Al₂Zr₃ for Al16. The phase prediction results suggest that the as-cast Al_x (x = 12, 14, 16) MEAs may consist of the BCC and HCP-structured phases, while Al_x (x = 18, 20) may be composed of HCP together with Zr-Al intermetallic.

The commonly used metastable phase prediction method and the d-electron alloy design method, developed based on the DV-Xα Cluster molecular orbital calculation, can describe the alloying effect successfully. Many Fe-based, Ni-based, and Co-based superalloys have been designed by the d-electron method, and it has also been widely used in medical titanium alloys with a low elasticity modulus. Morinaga et al. [24] drew a $\overline{Bo}-\overline{Md}$ empirical diagram based on d-electron theory, which can predict the phase stability of Ti alloys. $\overline{Bo}$ is the bond order and describes the bond strength between Ti and the alloying elements, where $\overline{Bo}={\sum }_i{x}_i{Bo}_i$; $\overline{Md}$ relates to the energy level of the d-orbitals and takes the atomic radius and the electronegativity of the elements into account, where $\overline{Md}={\sum }_i{x}_i{Md}_i$. In this diagram, different regions refer to HCP to BCC structure with α, α″, ω, or O phase were devised, and different deformation mechanisms such as dislocation slip and machine twinning were also marked. The Ms = RT is the line where the temperature of the martensitic transformation is room temperature, meaning that the alloys located nearby are metastable. The $\overline{Bo}-\overline{Md}$ diagram has been successfully used to explore the Ti-rich Ti₃₅Zr_27.5Hf_27.5Nb₅Ta₅ HEA [11]. Herein, the $\overline{Bo}-\overline{Md}$ criterion is used to estimate the phase stability of the presented Al_x MEAs. As can be seen in Figure 2, the Al_x MEAs are located in the α-stable region below the Ms = RT line, which is consistent with the predicted results of thermodynamic calculation to a certain extent.

3.2. Phase and Microstructure Evolution for Different Al Addition

The phase constitutions of the Al_x MEAs were characterized by XRD. The results shown in Figure 3 demonstrate that Al in the range of 12 to 20 at.% has a significant effect on the phase constitution. In Al12 and Al14 MEAs, only a single HCP structured phase is detected, and the lattice constant decreases with Al addition because of the shifting of diffraction peaks towards large 2θ. For Al16, the main phase is orthorhombic α″, and for Al18, the BCC phase appears and becomes the primary phase in Al20 alloy. For Al18 and Al20, the diffraction peaks at about 26° and 47° suggest the existence of another phase, which is speculated to be the B2 structure according to the published works for AlTiVCr [25], and Al_xHfNbTiZr [26]. The XRD results imply that Al acts as a β-stabilizer and leads to the transformation from α/α′ to α″ then to β. The actual phase composition is not consistent with CALPHAD and Bo-Md predicted results.

Figure 4 shows the microstructure of the Al_x MEAs characterized by the secondary electron mode of the scanning electron microscope. The alloys all consist of equiaxed grains, which are primitive β grains. For Al12 and Al14, martensitic laths of different sizes are inside the grains. In Al16, the number of laths on the grain boundary decreases and the sizes get bigger, in which the residue β phase appears. Increasing the Al content to 18 at.%, no laths were observed inside the grains.

To characterize the microstructure evolution caused by Al addition in further detail, TEM investigations were conducted on the Al12 and Al20 MEAs. The TEM results for Al12 in Figure 5 show that the sizes of the laths vary greatly, including smaller laths with a width of less than 100 nm and larger laths with a thickness of 500 nm. The EDS (Energy Dispersive Spectrometer, AZtec X-MaxN80, Oxford, UK) images in Figure 5a imply that the distribution of the elements is relatively uniform in different laths, which were confirmed to be in the HCP structure by SAED (selected area electron diffraction) for the thick laths shown in Figure 5b, which is consistent with the XRD results. However, a stripe-like internal structure is found in thick laths, which may be the Moire fringe implied by the HRTEM image in Figure 5c. The HRTEM indicates that the thick lath and the laths under it share the same [1$\overline{2}$1$\overline{3}$] zone axis with a misorientation of 33.4°, and they do not share the same crystal planes. The unique α phase may bring a strong fine-grain strengthening effect.

The phase composition of Al20 is more complex than Al12, as shown in Figure 6. The EDS results also show the uniform distribution of these four elements. The bright-field TEM image in Figure 6b displays three different phases dispersed on the matrix, including the slender-acerose phase, the elliptical black precipitate phase, and the elliptical white precipitate phase. The SAED of the marked region in Figure b1 shows that the matrix is in the BCC structure, and one of the precipitate phases is the metastable ω phase, which is common in traditional Ti-Nb- [27] and Ti-Mo-based alloys [28], and has been reported to play an important role in fine-scale α precipitation in these β-Ti alloys. However, it is noteworthy that the ω phase is usually suppressed by Al or Ta addition in conventional β titanium [29]. From the above results, it can be seen that Al did not inhibit the ω phase in these Ti-Zr-Hf-Al MEAs. The dark field image in Figure 6b3 shows that the volume fraction of the ω phase is large. The HRTEM image in Figure 6c1 and its FFT in Figure 6c2 also show that the coherency stresses arise at the β/ω boundary, and the β/ω interface acts as heterogeneous α nucleation sites, implied by the weak diffraction spots marked by the gray arrow. Except for the ω phase, another metastable phase α″ was detected, as shown in Figure 6b4, namely, the slender acerose phase with less than a 10 nm width, which was also verified by the XRD pattern in Figure 3. The HRTEM image for the matrix in Figure 6c3 and its FFT in Figure 6c4 along the [100] axis display the existence of superlattice reflections, which is concomitant with the B2 phase. However, there is no obvious interface between BCC and B2, implying a completely coherent relationship between them. The results further proved the fact that Al acts as a β-stabilizer in these Ti-Zr-Hf-Al medium-entropy alloys.

3.3. Mechanical Properties Variation Caused by Al Addition

To understand the effect of microstructure on mechanical properties caused by Al addition, tensile and compression tests were performed at room temperature. Figure 7 shows the tensile curves, corresponding work-hardening rate curves of the Al12 and Al20 MEAs, and the compression curve of Al20. The tensile properties, including yield strength σ_0.2, tensile strength σ_b, tensile fracture strain δ, and the elastic modulus E, are listed in

Table 2. Tensile properties including the yield strength, fracture strength, Young’s modulus, and fracture elongation of Al_x MEAs.

Alloys	Al12	Al14	Al16	Al18	Al20
σ0.2/MPa	1335	1796	-	-	721
σb/MPa	1528	1937	1731	1377	1346
E/GPa	115	102	966	821	678
δ/%	5.8	2.5	2.0	1.8	6.8

. It can be seen that Al shows an obvious effect on the mechanical properties. The Al12 and Al14 MEAs show extremely high strength comparable to high-strength steel, for σ_0.2 increasing from 1335 MPa to 1528 MPa and σ_b increasing from 1796 to 1937 MPa, with the Al content increasing from 12 to 14 at.%. The deformation capability deteriorates with Al; for Al12 and Al14, the fracture elongation is above 2.5%; however, Al16 and Al18 cannot yield before fracture with the fracture strengths of 1731 and 1377 MPa, respectively. The deformation behavior of Al20 shows significant differences compared with the other four alloys. After yielding at 721 MPa, an obvious strain-hardening stage appears, and the σ_b of 1347 MPa reaches a 6.8% fracture elongation.

The work-hardening rate curves of the Al12 and Al20 alloys are shown in Figure 7c,d. For Al12, the work-hardening rate decreases with strain and reduces to 0 at 3.5%, meaning that the true stress reaches its maximum at this strain, which is a typical phenomenon for α titanium. The strain-hardening rate of Al20 decreases first before 2.5% and then increases with strain, which is always above 5 GPa throughout the entire plastic deformation stage. The appearance of metastable α″ and β phases in Al20 suggests that the stress-induced martensitic transformation (SIMT) from the metastable β phase to the orthorhombic α″ phase may happen during deformation. The compression curve for Al20 exhibits double yielding, which is a typical symbol for the SIMT.

The SEM images of the fracture morphology of the Al_x MEAs are shown in Figure 8. For Al12 to Al18, the fracture surface shown in Figure 8a–d exhibits cleavage fracture, and the cracks cross the grain boundaries because of the strong intragranular solid solution strengthening and weak grain boundary. The Al20 alloy has two kinds of morphologies including the following: one is the pronounced intergranular fracture mode with plenty of vein-like patterns, as shown in Figure 8e, and the other is the almost paralleled slated fracture, which is the typical morphology of the SIMT.

4. Discussion

4.1. The Abnormal β-Stability Ability of Al in the Ti-Zr-Hf-Al System

In traditional Ti alloy, Al is an α-stabilizing element that increases the T_β (the temperature at which the α phase transfers to the β phase) and accelerates the formation of the α phase, which conforms to the calculated phase diagram showing that Al inhibited the precipitation of the β phase from the liquid phase. However, in fact, with an increasing Al content, the phase of Al_x MEAs transfers following α′(α)→α″→β + ω + B2. Specifically, the Al12 MEA is composed of the α phase, while the Al20 MEA mainly consists of the β phase, meaning that Al acts as a β stabilizer in Ti-Zr-Hf-Al MEAs, which goes against conventional wisdom. The reason may be that MEAs, by casting from liquid, are in the metastable state, and the metastable phase does not comply with the law of phase formation in a thermodynamic equilibrium state in the Thermo-calc calculated phase diagram.

The Bo-Md diagram in Figure 2 also cannot clearly explain the β phase stability of Ti-Zr-Hf-Al MEAs with a high alloying element content. The reasons include the following: (1) Bo describes the bond strength between Ti and the alloying elements. The “alloying elements” in the Ti-Zr-Hf-Al MEAs are more than that in traditional Ti-alloys. Neglecting the interaction among the “alloying elements” in MEAs will introduce errors, especially in the present Ti-Zr-Hf-Al system, where the enthalpy of mixing between Zr and Al is −44 kJ/mol, meaning that the interaction between these two elements is strong. (2) The values of Bo are different in different structures. For example, the Bo of Al in HCP-Ti is 3.297, but it is 2.426 in BCC-Ti. For Ti-Zr-Hf-Al MEAs with unknown structures, which values should be chosen is unclear. (3) MEAs with plenty of alloying elements are always located in the external region with few experiential cases. For the Ti₃₅Zr_27.5Hf_27.5Nb₅Ta₅ [11] and TiZrHfNb_x [30] alloys, they all lie neighboring the extension cord of Ms = RT, while there is no adequate experimental evidence to prove that the extension is reasonable.

The β-stabilizing capability of Al is different from traditional β stabilizers like Nb, Ta, or V, which decrease T_β. Al also displays the α-stabilizing effect in annealed alloys. The DSC curves of Al12 and Al20 are shown in Figure 9. Different from the simple phase transition behavior of Al12, there are two peaks in the Al20 alloy, namely, there is an exothermic peak from 597 to 708 °C, which responds to the decomposition reaction of metastable β and α″ to α and the stable β phase. The T_β values are 849 °C and 870 °C for Al12 and Al20, respectively, which indeed increase with Al. Similar results are reported for (Ti₆₀Zr₄₀)_100−xAl_x (x = 0, 5, 10, 15) alloys [30]. In the Ti-Zr-Hf-Al system, Hf is in the same family as Ti and Zr, so Al shows the same abnormal β-stabilizing ability.

The abnormal β-stabilizing effect of Al may be attributed to the high alloying in the Ti-Zr-Hf-Al system, like that in the Ti-Zr-Al system [23]. In these complex concentrated alloys, the heavy addition of Zr and Hf changes the movement of the outermost and subshell electrons, and the hybridization states of Ti and Al are altered. High cohesive energies could be caused by the alloying elements, which may be one reason for the abnormal β phase stability in Ti-Zr-Hf-Al alloys compared with low-alloy Ti alloys [23]. In addition, the local stress shows a positive correlation with mismatches among elements, and the addition of Al with a smaller atomic radius compared with Ti, Zr, and Hf results in extra local stress, postponing the long-range ordered martensitic transformation from β to α′. Instead, the transformation from the β phase to α″ martensite with a shorter distance of atom migration happened during casting. So, with the addition of Al, the major phase changes from α′ to α″ for Al12 to Al16. As the further increase in the Al content reaches 18 at.%, the transformation from the β phase to α″ is suppressed [23].

In Al20 MEAs with a large amount of Al content, there is lots of the ω phase, which violates the suppression effect of Al on the ω phase in the traditional metastable β titanium. The ordered B2 phase is often observed in Al-containing high-entropy alloys. With high electron density (the outermost shell holds three electrons) and a high Fermi level (small work function and high ionization tendency), Al prefers to transfer electrons to TMs, such as Ni, Co, Fe, Cr, and Ti, and forms intermetallic compounds with covalent bonds among the elements [31]. So, the large amount of Al addition into the Ti-Zr-Hf-Al system promoted the localized chemical order.

4.2. The Effect of Al on the Deformation Mechanism

Al acts as a β stabilizer, promoting the transformation from the HCP structure to the BCC structure, which finally changes the deformation behavior in these Ti-Zr-Hf-Al MEAs. It is well known that Al is the most important strengthening element in commercial titanium. For the Al12 MEA, the solid solution strengthening effect endows the alloy with high tensile strength up to 1528 MPa, which is higher than most of the FCC FeCoNiCrMn, BCC FeCoNiCrAl, and refractory HEAs in tensile properties. At the same time, it keeps a 5.8% elongation. The deformation mechanism changes from dislocation slip in the HCP crystal structure of Al12. As the Al content increased, the phase transformation from α to α′ and then to α″ happened, and a deformation mismatch between the hexagonal phase and highly solid solution strengthened, and the orthorhombic phase brought in the premature rupture of Al14 to Al18.

As the Al content increases up to 20 at.%, it acts as a β stabilizer, improving the formation of the metastable β and α″ phases, and the strong interaction between Al and other titanium group elements improves the formation of the ordered B2 phase. The phase combination makes the deformation behavior of Al20 complex, and high work-hardening happens under the extra stress. The phase evolution of Al20 before and after deformation was studied by using XRD, and obvious structure transformation occurred, as shown in Figure 10. After failure, the relative content of the B2 phase decreased and that of the BCC phase increased, which means that the order to disorder (B2 → BCC) transformation happened. The order/disorder transition under deformation was often observed in binary intermetallic compounds with the B2 or L1₂ structures. In addition, the SIMT from β to α″ occurs, which introduces an intensive strain-hardening effect by dynamic strain–stress partitioning between the β and α″ phases. It is common that the SIMT could also introduce good plastic deformation capacity, such as TiZrHfTa_x [12] and TiZrHfNb_x [30] HEAs. However, tensile elongation is limited to 6.8% for the Al20 MEA, which is due to the existence of the ω phase. The dispersive fine precipitation hinders the dislocation movement. The dislocation pile-up introduces high tensile strength but low deformation capacity. So, the phase transformation strengthening, including B2 to BCC and BCC to α″, together with the precipitation strengthening, brings about a high work-hardening ratio.

5. Conclusions

In this work, a series of (Ti₂ZrHf)_100−xAl_x (x = 12, 14, 16, 18, and 20, at.%, named Al_x) MEAs were prepared and characterized. The abnormal effects of Al on the phase, microstructure, and deformation behavior were investigated. The conclusions are summarized as follows.

1. Al acts as a β stabilizer in the Ti-Zr-Hf-Al system. As its content changes from 12 to 20 at.%, the phase evolution from α′(α)→α″→β + ω + B2 is not consistent with the α-stabilizing effect of Al in traditional titanium alloys. The phase transformation caused by Al cannot be predicted by the CALPHAD method or the Bo-Md diagram because of the complex interactions among composition elements.

2. The deformation behavior of Ti-Zr-Hf-Al MEAs changes with Al content. In the Al12 and Al14 MEAs, Al promotes solid solution strengthening, and the alloy with the HCP structure shows an extremely high strength up to 1937 MPa, which is rare in high-entropy alloys. But the deformation capability deteriorated with Al addition from 12 to 18 at.%.

3. In the Al20 MEA, Al not only acts as a weak β stabilizer, improving the formation of the metastable β and α″ phases, but it also promotes the localized chemical order, leading to the generating of B2. The phase transformation during deformation, including B2 to BCC and BCC to α″, together with the precipitation strengthening, brings about a high work-hardening ratio.