Realizing superior strength-ductility combination in dual-phase AlFeCoNiV high-entropy alloy through composition and microstructure design

Achieving high strength and ductility simultaneously in high-entropy alloys (HEAs) has always been a challenge, and recent studies indicate that duplex microstructure provides a strategy to solve this contradiction. We prepared Al0.25FeCoNiV HEA with L12/B2 fine-grain duplex microstructure and achieved a remarkable combination of the ultimate tensile strength (1530 MPa) and ductility (20%), which clearly outperforms other duplex HEAs. Owing to the formation of fine-grain microstructure and the hindrance of the B2 phase, high yield strength has been achieved. Furthermore, the excellent work-hardening ability of the L12 phase and the dispersion of the B2 phase ensured uniform deformation and high strength. GRAPHICAL ABSTRACT IMPACT STATEMENT Al0.25FeCoNiV HEAs exhibit a remarkable combination of ultimate tensile strength and ductility through L12/B2 fine-grain duplex microstructure, which clearly outperforms other duplex HEAs.


Introduction
High-entropy alloys (HEAs), a new type of alloys with more than four principal elements in equal or nearly equal atomic percentages, have attracted increasing attention owing to their unique microstructure and properties; for example, HEAs with body centered cubic (BCC) structure exhibit high strength and those with face centered cubic (FCC) structure exhibit high ductility [1][2][3]. However, nearly 10 years of research have shown that it is difficult to overcome the strength-ductility trade-off in single-phase HEAs [4]. Fortunately, the microstructures and mechanical properties of HEAs were found and can be tuned by modifying the chemical composition and thermal processes. Recent studies suggest that a dual-phase structure is effective in improving the mechanical properties  [5] demonstrated that the strength-ductility trade-off in metastable dual-phase Fe 50 Mn 30 Co 10 Cr 10 alloys can be resolved by varying the Mn content to generate transformation-induced plasticity (TRIP). Lu et al. [6] proposed a eutectic HEA with FCC/BCC lamellar structure, AlCoCrFeNi 2.1 , which combines the advantages of BCC and FCC phases. Furthermore, Shi et al. [7,8] obtained ultrafine-grain microstructure through heavy cryorolling and annealing, which exhibited ultimate tensile strength (UTS) of 1.6 GPa and ductility of 14%. However, the design of dual-phase alloy systems that exhibit TRIP or eutectic lamellar structure is still a challenge because the complex composition makes it difficult to predict the microstructure, often requiring several trial-and-error experiments.
Previous reports suggest that FCC/BCC dual-phase structure can be obtained by adding Al in FCCstructured HEAs, and the amount of BCC phase can be controlled within a certain range [9][10][11][12]. This is based on the research of the valence electron concentration (VEC) effect on the phase stability in solid solution forming HEAs, BCC phases are stabilized at a lower VEC (< 6.87), while FCC phases are stabilized at a higher VEC ( > 8.0), and in the intermediate VEC (6.87 < VEC < 8.0), both FCC and BCC phases exist [13]. Although the introduction of the BCC phase improves the mechanical properties, a low amount causes limited enhancement in the strength, and a high amount makes the alloy brittle. For example, Li et al. prepared Al 0.3 FeCoCrNiMn HEA with ultrafine microstructure through severe cold rolling and annealing treatment and then achieved yield strength (YS) of 1 GPa, but with low strain-hardening ability and plasticity less than 10% [14]. In addition, other hard and brittle phases, such as the σ phase, may form during the annealing process with the increase of Al contents [14]. Thus, it is difficult to achieve strength-ductility balance in HEAs through this strategy in current studies.
Vanadium (V) is an optimal element for strengthening FCC high-entropy alloys, and V is a key element in improving the Hall-Petch coefficient to enhance yield strength, such as CoNiV and V 10 Cr 15 Mn 5 Fe 35 Co 10 Ni 25 HEAs [15,16]; Meanwhile, V is easy to form intermetallic compounds and reducing the V content can effectively avoid the generation of intermetallic phases [17]. Hence, an equiatomic single-phase FeCoNiV HEA with FCC-L1 2 phase transition occurring at approximately 785°C was selected as the base alloy. The VEC of the Al element is very low, only 3. Therefore, the introduction of the Al element can reduce the VEC of the system and promote the transformation from FCC to BCC phase. In this paper, we introduced Al in FeCoNiV HEAs to control the VEC of the system and design an FCC/BCC duplex microstructure.

Materials and methods
The Al 0.25 FeCoNiV, referred to as Al025 was fabricated by arc melting under a high-purity argon atmosphere. The purities of all the elements are ≥ 99.95 wt. %. The liquid-state alloys were cast into a copper mold with dimensions of 7 × 10 × 100 mm 3 , samples taken from this state are denoted as 1#. The ingots were homogenized at 1100°C for 10 h in an argon-protected environment and cooled to room temperature with the furnace (2#). Then the ingots were cold-rolled at room temperature with a reduction ratio of 90% (thickness: 1 mm). Subsequent recrystallization and annealing treatments were performed in quartz ampules, which were then water quenched. Samples subjected to recrystallization at 900°C for 1.5 h are denoted as 3#. The samples subsequently annealed at 755°C for 10 min are named 4#.
The thermal properties of the HEAs were investigated by differential scanning calorimetry (DSC) with 96 Line equipment (SETARAM instrumentation, France) operated in a N 2 atmosphere from 50 to 1100°C at a rate of 10°C/min. The different phases in the HEAs were determined by X-ray diffraction (XRD; D8 ADVANCE Da Vinci, Bruker, Germany) and Synchrotron radiation X-ray diffraction (SXRD; BL14B1 beamline of Shanghai Synchrotron Radiation Facility, China). The phase characterization and deformation behavior were investigated using energy disperse spectroscopy (EDS), electron backscatter diffraction (EBSD), and transmission electron microscopy (TEM). The EBSD and EDS investigations were performed in a scanning electron microscope (SEM; LYRA3 GMU, TESCAN, Czech Republic). The TEM and EDS investigations were performed using a JEM-2100F (JEOL, Japan) instrument at an accelerating voltage of 200 kV. The samples for TEM analysis were prepared by conventional electropolishing of foils (mechanically pre-thinned to 60 μm) in a mixture of 10% HClO 4 and 90% C 2 H 5 OH at 25 V and −25°C. Tensile tests were performed using a universal testing machine (MTEST 5000 W Tensile Stage, GATAN, UK) at a quasistatic strain rate of 0.1 mm/min and room temperature. Plate-type tensile specimens (gauge length: 4 mm, crosssection: 1 × 2 mm 2 ) fabricated by electrical discharge machining were used for the tensile tests. The strain was measured synchronously during the tensile process by digital image correlation using Vic-2D software (Correlated Solutions, USA). For this, a speckle pattern was applied on the surface of the tensile specimens.

Results and discussion
The DSC curves of the as-cast HEAs in Figure 1(a) show that the exothermic and endothermic peaks in the heating and cooling processes of Al025 HEA, respectively. And this phenomenon corresponds to the solidstate phase transformation of the studied alloy. The XRD results of the as-cast HEAs confirm that the addition of Al promotes the formation of the BCC phase, and Al025 exhibits the FCC/BCC dual-phase structure as shown in Figure 1(b). And no phase transition can be seen between as-cast (1#) and homogenized state (2#). Detail investigation on the microstructure including EBSD and TEM analysis was carried out and 1# Al025 HEA exhibits a dual-phase microstructure consisting of FCC and BCC phases; the fraction of the BCC phase is 12 vol.%, as shown in Figure 1(c-f). The BCC phase is distributed at the grain boundaries and the interdendritic regions of the FCC phase, and there seems to be an orientation relationship between the BCC phase and the FCC phase. The V content (28.80%) in the BCC phase is significantly higher than that in the FCC phase (21.80%). After homogenization treatment, Al025 still exhibits duplex microstructure and component dendrites were eliminated. The Al content (13.50%) in the BCC phase is evidently higher than that in the FCC phase (4.30%) (Figure 1(g-i)). The detailed chemical compositions of the HEAs are presented in Table 1. It is important that the FCC and BCC phases in the 2# sample are found to translate to chemical ordered L1 2 and B2 phases, respectively, from the TEM studies (Figure 1(j-l)). And the dark-field (DF) image corresponding to the superlattice spot (100) of the L1 2 phase shows clear anti-phase boundaries (APBs) and domain structure. The orientation relationship is not affected, and the orientation of the partial two phases still satisfies the K-S relationship ({111} L12 //{110} B2 ; 110 L12 // 111 B2 ). Figure 2 shows the detailed microstructures of the recrystallized (3#) and subsequent annealed (4#) Al025 HEAs. The SXRD results (Figure 2(a)) show that after cold-rolling and recrystallization treatments, Al025 alloy exhibited the FCC/BCC dual-phase structure, annealed Al025 alloy was composed of L1 2 phase and B2 phase, and no other intermetallic compounds were observed. The 3# Al025 HEA exhibits a fine-grain dual-phase microstructure and similar element distribution inhomogeneous, as shown in Figure 2(b-d). The average grain size of the L1 2 phase is approximately 2 μm, while that of the B2 phase is approximately 680 nm. The phase distribution map (Figure 2(c)) shows that there are two types of B2 phases: one consisting of large clusters (approximately 1 μm), and the other consisting of small grains (approximately 200 nm) dispersed on the L1 2 grain boundaries and inside the L1 2 grain. The EDS maps in Figure 2(d) show that the Al element distributes differently between B2 and L1 2 phases, whereas the distribution of V in the two phases is relatively uniform. In addition, the two types of B2 phases have similar compositions. When the recrystallization temperature is decreased to 800°C, it is found that the L1 2 phase is partially recrystallized while the B2 phase is completely recrystallized, as shown in Figure 2(e). In particular, it is observed that there is a large number of recrystallized B2 grains in the nonrecrystallized L1 2 region. This confirms that the B2 phase is more likely to nucleate and recrystallize in the recrystallization process, and it seems difficult to grow up, which then can pin L1 2 grain boundaries and refine the L1 2 grains. The EBSD and TEM results indicate a dual-phase microstructure consisting of FCC and BCC phases and there is no orientation relationship, especially between the BCC phase and the FCC phase ( Figure  2(b,f)). After annealing at 755°C for 10 min, the 4# Al025 HEA exhibits a fine-grain dual-phase microstructure consisting of L1 2 and B2 phases, as shown in Figure  2(g). And the dark-field (DF) image corresponding to the superlattice spot exhibits irregular ordered domain structure, and combined with TEM-EDS, a single L1 2 phase was confirmed in this region (Figure 2(h,i)). B2 phase was also confirmed in this way and TEM-EDS reveals its characteristics of rich Al and poor Ni elements ( Figure  2(j-l)). Figure 3 shows the tensile properties of these HEAs at room temperature. The YS of fine grains 3# and 4# Al025 HEA (1100 MPa) is significantly higher than that of coarse grains 1# and 2# Al025 HEA; it is even higher than the ultimate tensile strength of 1# HEA, as shown in Table 2. In contrast to 3#, 4# Al025 HEA exhibits little improvement in the YS owning to the    [7]. it is evident that this Al025 HEA exhibits the best mechanical properties [6,7,14,[18][19][20][21]. To elucidate the strengthening and deformation mechanisms of Al025 HEA, the work-hardening rate was analyzed as a function of the plastic strain, as shown in Figure 3(c). First, by comparing the workhardening curves of 2# and 4#, grain size and B2 phase distribution have an influence on the early stage of the strain-hardening process and have little effect on the subsequent strain-hardening behavior. Second, the excellent work-hardening ability in 2# and 4# Al025 HEAs should be attributed to the ordered structure of the L1 2 and B2 phase by contrast with the disordered structure of Al025 HEAs. According to the previous studies on ordered alloys [22][23][24][25][26], both the L1 2 phase and B2 phase have been reported to improve work-hardening rate, which is related to the superdislocations configuration and slip behavior. Third, because of the previous synthesis effect, a homogeneous deformation can be maintained and good ductility is expected in the Al025 alloy. The representative fracture surfaces of 4# Al025 HEAs are shown in Figure 3(d). The fracture is almost ductile dimple fracture, which is due to the considerable deformation capacity of the L1 2 phase [27]. The fracture morphology also includes intergranular facets marked by white dotted lines, with a size of less than 3 μm, which is consistent with the grain size of the L1 2 phase. In addition, there are dark holes near the intergranular facets, and the size is consistent with B2 on the grain boundary.
To further investigate the deformation behavior in the fine-grain dual-phase HEA, detailed characterizations were carried out on the 4# Al025 HEA sample. Figure  4(a)-(c) shows the EBSD inverse pole figure, phase map, and kernel average misorientation (KAM) distribution map of the 4# Al025 HEA subjected to a strain of 12% at room temperature. According to the definition of KAM, a low KAM value indicates homogeneous plastic deformation among the local regions. The L1 2 phase bears the main strain, while the strain gradient is mainly distributed within the B2 phase and the interface between L1 2 and B2 phases. Predictably, the stress accumulated in the phase boundaries between the B2 phase and L1 2 phase cannot be released because of the limited deformation ability, and cracks can easily initiate along the phase boundaries. Further TEM investigations were shown in Figure 4(d-f) to reveal the deformation mechanism in B2 and L1 2 phases. Figure 4(d) shows the BF image and SAED along the L1 2 [011] zone axis of 4# Al025 HEA. A large number of dislocations were found distributed in the L1 2 phase and little dislocations can be seen in the B2 phase. SAED results indicate the existence of stacking faults (SFs) and lots of SFs can be seen in the L1 2 phase, as shown in Figure 4(e). The deformation ability of the L1 2 phase in Al025 HEA is stronger than that of traditional L1 2 phases, such as Ni 3 Al and Ni 3 Si; and the current widely accepted explanation of strain hardening is that the special jogs generated by dislocation intersection or cross-slip bring additional hardening reported in Cu 3 Au and (Fe, Co, Ni) 3 V [23,24,28,29]. According to the previous studies and the current results [27], we believe that L1 2 phase phases have a dominant role in the straightforward hardening and the deformation mechanism of the L1 2 phase in Al025 HEA needs to be studied further. The present work mainly focuses on the deformation ability of the B2 phase and the interface between the L1 2 and B2 phases, since the introduction of the B2 phase in AlFeCoNiV HEA affects the plasticity of the material. As shown in Figure 4(f), abundant dislocation pile-up at the interface between the L1 2 and B2 phases, reflects the influence of the B2 phase on the strength and plasticity.
The evolution of the dual-phase microstructure in Al025 during heat treatment and deformation is summarized in Figures 4(g-i), which explains the origin of strength and plasticity. The grains of the L1 2 phase is relatively large in the homogenized sample (2#), and the B2 phase is distributed at the grain boundaries and inside the grains. The BCC phase is crushed by cold-rolling deformation, and the nucleation of the BCC phase is superior to that of the FCC phase in the subsequent recrystallization process. This results in microstructure refinement as shown in Figure 4(g,h). The hindrance of the B2 phase on the dislocation slip also contributes to the YS. Moreover, since V can significantly improve the Hall-Petch coefficient, the pinning of L1 2 grain boundaries by the BCC phase to obtain fine grains can significantly improve the YS of Al025 alloys, which is different from previous dual-phase HEAs. The L1 2 phase plays an important role in the strain-hardening process, and deformation mainly occurs in the L1 2 phase, as shown in Figure 4(i). However, as deformation progresses, the deformation of the phase boundaries and that of the grain boundaries between B2 grains are not synchronized owing to the poor deformation ability of the B2 phase. This results in an intergranular fracture which is confirmed by the fracture surface. Increasing the deformation of the BCC phase during the cold rolling process to obtain a more dispersed BCC phase is beneficial for further improving the mechanical properties of these AlFeCoNiV HEAs.
In summary, high strength-ductility combination was achieved in Al025 HEA with an L1 2 /B2 fine-grain duplex microstructure via the coordination of L1 2 /B2 in the deformation process. The L1 2 phase is a prerequisite for plastic deformation, and fine grains significantly contribute to the YS owing to the high Hall-Petch coefficient. Moreover, the unique strain-hardening behavior of the L1 2 phase ensures uniform deformation of the material, and the dispersed B2 phase prevents direct expansion of the crack. The fine B2 phase (even though in small amounts) not only refines the L1 2 grains but also contributes to the YS. The role of the B2 phase in Al025 HEA is different from that in previously reported duplex-microstructure alloys. This finding is beneficial for designing a strategy to develop new duplex HEAs with outstanding strength-ductility combination.