Microstructural origins of high strength and high ductility in an AlCoCrFeNi 2.1 eutectic high-entropy alloy

Recent studies indicate that eutectic high-entropy alloys can simultaneously possess high strength and high ductility, which have potential applications in industrial fields. Nevertheless, microstructural origins of the excellent strength–ductility combination remain unclear. In this study, an AlCoCrFeNi 2.1 eutectic high-entropy alloy was prepared with face-centered cubic (FCC)(L1 2 )/ body-centered-cubic (BCC)(B2) modulated lamellar structures and a remarkable combination of ultimate tensile strength (1351 MPa) and ductility (15.4%) using the classical casting technique. Post-deformation transmission electron microscopy revealed that the FCC(L1 2 ) phase was deformed in a matter of planar dislocation slip, with a slip system of {111} <110>, and stacking faults due to low stacking fault energy. Due to extreme solute drag, high densities of dislocations are distributed homogeneously at {111} slip plane. In the BCC(B2) phase, some dislocations exist on two {110} slip bands. The atom probe tomography analysis revealed a high density of Cr-enriched nano-precipitates, which strengthened the BCC(B2) phase by Orowan mechanisms. Fracture surface observation revealed a ductile fracture in the FCC(L1 2 ) phase and a brittle-like fracture in the BCC(B2) lamella. The underlying mechanism for the high strength and high ductility of AlCoCrFeNi 2.1 eutectic high-entropy alloy was finally analyzed based on the coupling between the ductile FCC(L1 2 ) and brittle BCC(B2) phases.


Introduction
High-entropy alloys (HEAs), which emerged in 2004, are solid solution multicomponent alloys that contain more than four principal elements in equal or nonequal atomic percentage [1][2][3][4][5]. HEAs are a new kind of alloys because they are different from conventional alloys that have one or two principal elements as main components.
The review of literature indicates that two kinds of HEAs have been widely investigated in the last 10 years: single-phase FCC HEAs and BCC HEAs. Single-phase FCC HEAs generally have high tensile ductility but low yield strength. Examples are CrMnFeCoNi alloy with an elongation to failure of ~50% and yield strength of ~410 MPa [7], Fe 40 Mn 40 Co 10 Cr 10 alloy with a tensile ductility of ~58% and low yield strength of ~240 MPa [17], and Fe 40 Mn 26 Ni 27 Co 5 Cr 2 alloy with a total elongation to failure of ~58% and low yield strength of ~95 MPa at room temperature [18] level (about 4-4.5 GPa) but low plasticity (less than 0.2%) [19]. Therefore, for the singlephase HEAs, it seems difficult to achieve a balance between high strength and high ductility. In addition, the industrial applications of HEAs are restricted certainly by poor castability, liquidity, and composition segregation.
Recently Lu et al. [20] proposed a designation concept of eutectic high-entropy alloys (EHEAs) to combine the high strength of BCC HEAs and high ductility of FCC HEAs and produced an AlCoCrFeNi 2.1 EHEA with regular FCC(L1 2 )/BCC(B2) lamellar structures and an excellent combination of high strength and high ductility. More recently, the mechanical properties of AlCoCrFeNi 2.1 EHEA were optimized by thermal mechanical processes, that is, muti-pass cold rolling to a 90% reduction in thickness and subsequently annealed at 800-1200°C for 1 h [21,22]. As a result, the processed HEAs had tensile ductility more than 10% and high tensile strength greater than 1 GPa. This indicated that the properties of the EHEA could be successfully tailored using thermo-mechanical processing for a wide range of engineering applications. However, the microstructural origins of excellent strength-ductility combination and underlying deformation mechanisms of EHEAs still need systematic investigations.
In this study, transmission electron microscopy (TEM) and atom probe tomography (APT) were used to systematically characterize the microstructural characteristics of post-deformed AlCoCrFeNi 2.1 EHEA to find answers to the aforementioned problems.

Experiments
The master alloy of the eutectic AlCoCrFeNi 2.1 (elements in atomic ratios) was prepared from commercially pure elements (99.9 wt.% for Al, Co, and Ni; 99. 5-99.6 wt.% for Cr and Fe) in a ZrO 2 crucible of a vacuum induction melting furnace. The ZrO 2 crucible was first heated to 600°C and held for 1 h to remove the water vapor. After the elements (approximately 2.5 kg) were put into the furnace, the furnace chamber was evacuated to 0.06 Pa and backfilled with high-purity argon gas to reach 0.06 MPa. The elements were finally melted, superheated, and poured into a high-purity graphite crucible with an inner length of 220 mm, upper inner diameter of 62 mm, and bottom inner diameter of 50 mm. The pouring temperature was set to be 1500°C. A TRTM-2CK infrared pyrometer was used to monitor the temperature with an absolute accuracy of ± 2°C. A Walter + bai LFM 20kN tensile testing machine was used for tensile testing at room temperature with a normal strain rate of 1 × 10 -3 s -1 . The flat dog bone-shaped tensile samples had a gauge dimension of 20 × 3 × 2 mm 3 . The strain was measured using a standard non-contacting video extensometer. Three tensile specimens were measured to obtain reliable results.
Microstructure and composition analyses were carried out by means of x-ray diffraction (XRD), electron back-scattered diffraction (EBSD), TEM, and APT.
Specifically, the XRD analyses of crystalline structures were performed on a Bruker D8 with Cu radiation target scanning 2θ from 20 to 80 degrees. EBSD of as-cast specimen was conducted using a high-resolution field emission Carl Zeiss-Auriga-45-66 scanning electron microscope (SEM) equipped with a fully automatic Oxford Instruments Aztec 2.0 EBSD system (channel 5 software). Before EBSD, the specimens were mechanically polished and then electro-polished in an electrolyte containing 90 vol.% acetic acid and 10 vol.% perchloric acid using a voltage of 50 V and polishing time of 45 s in Buehler electromet-4. TEM observations were conducted in a FEI-Tecnai G 2 20 S-TWIN microscope operated at 200 kV, and the high-resolution TEM (HRTEM) was conducted on a Titan G2 60-300. TEM specimens were prepared as follows: the tensile deformed gauge parts were carefully ground to foils with a thickness of about 50 μm, punched into disks with a diameter of 3 mm, and finally electro-polished to an electron-transparent thickness in an aqueous electrolyte containing 10% perchloric acid and 90% ethanol at -25°C using a twin-jet polishing system. The HRTEM specimen was prepared by means of ion milling on a GATAN 691 to avoid the different electro-polishing rates of the two eutectic phases.
The APT specimen was prepared by electro-polishing combined with a focus iron beam (FIB) cutting along the phase interface. A low-energy (5 keV) Ga beam was used for final ion milling to minimize beam damage. The APT experiment was performed with a local electrode atom probe (LEAP 4000X Si) under ultraviolet laser pulsing at a laser energy of 40 pJ, a pulse repetition rate of 200 kHz, and a target evaporation rate of 0.5% per pulse at 25 K. APT data reconstruction and quantitative analysis were performed using a CAMECA visualization and analysis software (IVAS) 3.6.8. Figure 1a shows the XRD pattern of the as-cast AlCoCrFeNi 2.1 alloy. The as-cast AlCoCrFeNi 2.1 alloy consisted of FCC(L1 2 ) and BCC(B2) dual-phase. The L1 2 and B2 structures, which were further confirmed by subsequent TEM-SAEDs (select area electron-beam diffraction patterns), were consistent with the results reported in the literature [21,22]. Figure 1b shows the large-area EBSD phase mapping of typical FCC(L1 2 )/BCC(B2) lamellar microstructures. Fine BCC(B2) lamellae (about 2 μm thick, yellow color) were parallel to each other and distributed in the FCC(L1 2 ) phase (cyan color). In addition, some coarse BCC(B2) lamellae and islands also appeared. The volume fraction of FCC(L1 2 ) and BCC(B2) phases was about 66.3% and 31.2%, respectively. Moreover, some amounts of high-angle grain boundaries (marked in red) were formed during casting.

Initial structures and tensile testing results
Tensile testing showed that the as-cast AlCoCrFeNi 2.1 alloy possessed an excellent combination of high strength and high ductility owing to the uniform FCC(L1 2 )/BCC(B2) lamellar microstructures. Figure 2 shows the tensile engineering and true stress-strain curves of the as-cast AlCoCrFeNi 2.1 alloy. It is evident that the ultimate tensile stress was 1100 50 MPa and the ductility was 18 2%. When the engineering stress-strain ± ± curve (blue line) was converted into a true stress-strain curve (red line), the ultimate tensile stress and ductility were 1351 MPa and 15.4%, respectively .

Deformation of FCC(L1 2 ) phase
TEM observations were performed to reveal the deformation mechanisms and understand the mechanical properties. Figure 3a shows a representative bright-field TEM micrograph of the tensile tested AlCoCrFeNi 2.1 EHEA. Distinct modulated FCC(L1 2 ) and BCC(B2) lamellar structures were obviously observed based on different contrasts (bright and dark, respectively) in the TEM image. This was because the FCC(L1 2 ) phase was thinner than the BCC(B2) phase after twin-jet polishing. This chemical polishing result also suggested that the corrosion resistance of BCC(B2) phase was better than that of FCC(L1 2 ) phase. Figure 3b and 3c show the TEM-SAEDs of FCC(L1 2 ) and BCC(B2) phases, respectively. The presence of super-lattice spots obtained from the two phases (marked by cyan and yellow circles, respectively) revealed both ordered L1 2 ( Fig. 3b) and B2 (Fig. 3c) phases. The formation of ordered L1 2 structure is rarely reported in eutectic alloys in the literature, whereas the formation of ordered B2 phase is frequently reported owing to the strong negative enthalpy in the Al-Ni system [23][24][25]. The compositions of the present B2 phase were revealed as ~38 at.% Ni and ~37 at.% Al by the following APT analysis (see Table 1), which hinted that Ni and Al atoms alternately occupied the lattice sites in the B2 unit cell, and other element atoms were substitutional solution atoms [26][27][28].
The deformation mechanisms of FCC(L1 2 ) and BCC(B2) phases were different from each other. Figure 3d is a magnified TEM image of FCC(L1 2 ) phase region marked "A" in Fig. 3a. Several arrays of parallel mobile dislocations could be discerned, as pointed by the white arrows. The specimen tilting technique [29] was used to determine the slip plane of parallel dislocations, and as a result, the dislocation glide was clearly localized on a distinct set of {111} type FCC lattice planes. From the density of parallel dislocations in the array, one could judge that the dislocations were emitted from the FCC(L1 2 )/BCC(B2) phase boundaries. Besides high-density parallel dislocations in the FCC(L1 2 ) phases, the presence of some stacking faults (SFs, marked by black arrows in  Figure 4 shows bright-field TEM images of the tensile tested EHEA specimen from an alternative region. The BCC(B2) phase had a medium density of dislocations after tension deformation and high density of nano-precipitates ( Fig. 4c and 4d) (composition and structure characterized in the following section). The detailed analysis indicated that the dislocations in the B2 phase could be observed at the [001] zone axis and were mostly located in two slip bands, as displayed more clearly in Fig. 5a. These dislocations were identified in ( 10) and (110)  dislocations in B2 phases were also found (Fig. 5b). The arrow-shaped dislocations were formed in the following manner: when two straight parallel dislocations were pinned by a nano-precipitate on the tip, the contraction acquired an arrow shape. This result indicated that the precipitate-strengthening effect played an important role in the high strength of B2 phase.

Deformation of BCC(B2) phase
Briefly, the FCC(L1 2 ) phase deformed via dislocation planar slip and SFs to act as a soft phase, whereas the B2 phase acted as a hard phase reinforced by nano-precipitates.
The modulated hard and soft fine, regular eutectic lamellar structures of AlCoCrFeNi 2.1 EHEA created the excellent combination of high strength and high ductility.

Nano-precipitates by APT characterization
The tip specimen was cut by FIB and analyzed by APT analysis to discover the composition of nano-precipitates in the B2 matrix phase. The 3-D atom maps are shown in Fig. 6a Fig. 6a also reveals substantial spherical nano-precipitates enriching Cr atoms with sizes less than 20 nm in the B2 phase. Figure 6b further shows a size distribution of Cr-enriched nanoprecipitates. The proportion of Cr-enriched nano-precipitates with sizes less than 5 nm was more than 60% and that of nano-precipitates with sizes bigger than 10 nm was 11% (Fig. 6c).
The quantitative composition analyses of nano-precipitates and BCC as well as FCC phases were further conducted. Figure 7a is the 1-D proxigram of nano-precipitates showing elemental concentration profiles along the radial direction of their iso-surfaces.
The position at X = 0 was the interface (defined by the isosurface) of nano-precipitates and B2 matrix, whereas the positions at X > 0 and X < 0 were nano-precipitate interiors  AlCoCrCuFeNi [35] HEAs, which originated from the spinodal decomposition of compositional modulation [22]. However, the Cr-enriched nano-precipitates with the special structure inside the B2 phase were first detected in the AlCoCrFeNi 2.1 EHEAs.

SFs and phase boundary by HRTEM characterization
The qualitative characterization of SFs and atomic structures of eutectic dual phases was performed by HRTEM. Figure 9a  shown in Fig. 9c. These results indicated that the SF was another important deformation mode besides planar-slip aforementioned in the FCC(L1 2 ) phase.
HRTEM also uncovered the semi-coherent phase boundary between FCC(L1 2 ) and BCC(B2) phases (Fig. 9c). Therefore, the combination of eutectic dual phases was at the atomic level and strong enough to bear high stress. The  phase. The load was then imposed to the FCC(L1 2 ) phase, and due to the large strain hardening capability, the FCC(L1 2 ) phase was subjected to a perfect plastic deformation by necking into sharp lines (bright lines) without any dimples.

Discussion
Compared with the traditional metal materials, HEAs possess a specific atomic structure, which is the result of the mixing of elements with different atomic sizes, hence producing deformed lattices. The massive different atoms improve the solid solution strengthening and change the SFE. On the contrary, the frictional stress is assumed to be high due to the large lattice distortion caused by different atomic radii and moduli [19].
Therefore, the significant lattice distortion results in the high ductility and strain hardening of FCC HEAs in the absence of intermetallic phases. The lattice distortion might produce specific deformations of HEAs, which further determine their specific mechanical properties. Careful inspection of the literature, however, indicates that analyses of the deformation mechanisms of such complex alloys and especially the nature of defects are limited [9,11,30,[38][39][40].
The early study on the deformation mechanisms of FCC FeCrNiCoMn using TEM indicated that at a large range of temperatures (77-873 K), plastic deformation occurred by a planar dislocation glide on the classical FCC slip system, {111}<110> and SFs formed through dissociation of the unit dislocation b = a/2<110> in {111} plane into 1/6<112> Shockley partials [30]. When the plastic strain was larger than 20%, dislocations re-organized into cell structure, with the size range of 200-300 nm similar to other FCC alloys with low-to-medium SFE, at room temperature, and deformation twinning appeared at 77 K. The authors also suggested that the planar glide might be caused by the low SFE of the alloy or a short-range order effect. The planar dislocation glide was frequently reported in low SFE binary alloys, such as Cu-Al [36] and Cu-Zn [37]. A more recent study of the same FCC FeCrNiCoMn by means of in situ TEM observation revealed some unique deformation mechanisms. The planar slip was not continuous with extremely low velocity, which formed close-packed dislocation arrays and pile-ups in the activated {111} due to the strong drag effect of solute atoms [9]. In For the BCC HEA, an early study of TiZrHfNbTa HEA indicated that at the mesoscopic level, deformation was not homogeneous but often located in intragranular shear bands [38]. Deformation twins were also observed near grain boundaries. A thermally activated dislocation glide was suggested as the deformation mechanism by the authors based on their analyses of the strain rate and temperature dependences of flow stress [38].
Moreover, the authors also proposed a mechanism based on solid solution strain hardening to explain the deformation behavior of this alloy [11]. More recent investigations of the same BCC TiZrHfNbTa HEA by means of TEM revealed that the dislocation glide was controlled by the movement of screw dislocation at the beginning of deformation, and the deformation was rapidly localized in two distinct bands where dislocation dipoles, loops, and tangles were induced at a higher plastic strain [39,40]. All Bulk coarse-grained B2 NiAl single-phase alloys are brittle at room temperature.
This is because coarse-grained NiAl alloys possess only three independent slip systems that cannot produce cooperative deformation among grains [41]. One could speculate that the B2 rich-NiAl phase in the AlCoCrFeNi 2.1 EHEA with solid solution strengthening and Cr-enriched precipitate strengthening might be even more difficult to be plastically deformed. However, the results of the present study indicated no cracks at the phase boundary of the tensile deformed specimens (Fig. 3, 4, and 9), which suggested that both the FCC(L1 2 ) and BCC(B2) phases were deformed synchronously to about 18% elongation. The observed medium density of dislocations also provided enough evidence for the plastic deformation of BCC(B2) phase. The possible reasons for the enhanced plasticity of brittle BCC(B2) phase in the EHEA alloys might be the complex 3D backstress states subjected to the FCC(L1 2 ) phase. In the literature, the back stress has been verified to improve the plasticity of nanostructured surface layer when adjacent to a coarse-grained core [42]. The back stress was produced in the heterogeneous systems during deformation and maintained the deformation synchronously. In the AlCoCrFeNi 2.1 EHEA, the boundaries of hard B2 phase and soft FCC(L1 2 ) phase were semi-coherent (see Fig. 9c), and these were enveloped by the soft FCC(L1 2 ) phase. These conditions favored the occurrence of back-stress effect that could result in massive dislocation in the FCC(L1 2 ) phase accumulated at the phase boundaries and activate more dislocations in the BCC(B2) phase compared with the independent coarse-grained NiAl alloys in the uniaxial stressed state during a tensile test.

Conclusions
In this study, the AlCoCrFeNi 2.1 EHEA was prepared using the casting technique.