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Article

Load Partitioning and Strain Compatibility in a Non-Equiatomic Dual-Phase AlCoCrFeNi High-Entropy Alloy Processed by Forging

by
Pablo Pérez
1,*,
Sergio Perosanz
2,
Judit Medina
1,
Edurne Laurín
1,
Alberto Orozco-Caballero
1,
Rebeca Hernández
3,
Andreas Stark
4,
Norbert Schell
4,
Paloma Adeva
1 and
Gerardo Garces
1
1
Departamento de Metalurgia Física, Centro Nacional de Investigaciones Metalúrgicas (CENIM-CSIC), Avda. Gregorio del Amo 8, 28040 Madrid, Spain
2
E.T.S. de Ingeniería Aeronáutica y del Espacio, Universidad Politécnica de Madrid, Pza. Cardenal Cisneros 3, 28040 Madrid, Spain
3
División de Materiales de Interés Energético, Centro de Investigaciones Energéticas, Medioambientales y Tecnológicas (CIEMAT), Avd. Complutense 40, 28040 Madrid, Spain
4
Institute of Materials Research, Helmholtz-Zentrum Geesthacht, Max-Planck-Str. 1, 21502 Geesthacht, Germany
*
Author to whom correspondence should be addressed.
Metals 2026, 16(3), 300; https://doi.org/10.3390/met16030300
Submission received: 30 January 2026 / Revised: 1 March 2026 / Accepted: 4 March 2026 / Published: 8 March 2026
(This article belongs to the Special Issue Advanced Manufacturing and Processing Technology for Metallic Materials)
Browse Figures
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Abstract

The tensile and compressive behavior of hot-forged Al5Co35Cr30Fe20Ni5 high-entropy alloy (HEA) has been studied at room temperature. The forged HEA has a dual-phase microstructure consisting of a predominant face-centered cubic (FCC) matrix and a body-centered cubic (BCC) phase. The BCC phase embeds a low volume fraction of ordered BCC nanoparticles (B2 structure). During forging, the BCC phase recrystallizes more easily than the FCC phase. Yielding is controlled by the deformation of the FCC phase, although BCC grains assume an additional part of the load transferred by FCC grains, even during the elastic regime. During the onset of plastic deformation, slip is activated preferentially in the FCC phase in those grains that are favorably oriented for slip in planes (111). Dislocation pile-ups at FCC/BCC interfaces induce dislocation slip in the BCC phase. In the BCC phase, B2 particles act as effective obstacles to dislocation motion through the Orowan mechanism. As the deformation proceeds, dislocation activity causes an increase in the misorientation in both phases, resulting in the formation of subgrains whose boundaries are effective for blocking dislocation motion. The combination of high strength and ductility arises from the dual-phase FCC–BCC microstructure of the alloy. The load borne by the BCC phase partially relieves the stress applied to the FCC matrix, enabling the latter to continue deforming.

1. Introduction

High-entropy alloys (HEAs) have emerged in recent years as promising substitutes for superalloys in structural applications, particularly those requiring stability in extreme high-temperature environments [1,2]. The HEA concept was initially defined by equiatomic compositions designed to maximize configurational entropy and stabilize single-phase solid solutions [3,4]. HEAs exhibit four “core effects” that contribute to their promising characteristics: high configurational entropy, sluggish diffusion, severe lattice distortion, and the “cocktail effect”. High configurational entropy lowers the Gibbs free energy, thereby promoting the formation of solid solution phases, especially at elevated temperatures [5,6,7,8]. Sluggish diffusion kinetics contribute to enhanced thermal stability and resistance to grain growth [9]. Furthermore, severe lattice distortion, arising from the atomic size mismatch among multiple principal elements, significantly influences dislocation motion and provides substantial solid solution strengthening [10,11].
The solid-solution stability of many nominally single-phase HEAs is compromised during exposure to intermediate temperatures, and therefore, the definition has progressively broadened. Contemporary classifications now encompass alloys containing at least five principal elements with atomic concentrations between 5% and 35%, allowing for a wider compositional design space and improved phase stability [5,6]. Importantly, this conceptual expansion is not limited to composition but also includes microstructural design, where multiphase HEAs are intentionally developed as a strategy to tailor mechanical performance [5,6,7]. In particular, dual-phase FCC–BCC or FCC–HCP HEAs have attracted increasing attention due to their ability to combine the high ductility of FCC phases with the high strength of BCC/HCP phases, enabling synergistic mechanical behavior through phase-specific plastic deformation and load partitioning mechanisms [12,13,14,15,16,17,18].
Among the various HEA systems, the AlCoCrFeNi system—particularly near-equiatomic compositions—has been extensively studied as a leading candidate for demanding engineering applications [19,20]. Equiatomic AlCoCrFeNi HEAs typically exhibit a microstructure consisting of a mixture of body-centered cubic (BCC) and ordered BCC (B2) phases [3,6,19]. In contrast, non-equiatomic alloys display face-centered cubic (FCC) phases or complex multi-phase hierarchies depending on specific elemental ratios. For instance, increasing the aluminum content is known to promote BCC phase transition [21]. Conversely, deviations from equiatomic ratios can induce the precipitation of intermetallic phases, such as L12 or sigma phases, which significantly alter the mechanical properties [22,23]. Furthermore, the presence of ordered phases or nanoprecipitates in non-equiatomic alloys influences dislocation motion through mechanisms such as Orowan looping or particle shearing, thereby enhancing the overall strength [11,24]. Recent research has also highlighted the critical role of chemical short-range order (SRO) on defect behavior [25]. The presence of ordered atomic clusters within the nominally random solid solution alters the energy landscape for dislocation movement, thereby influencing the alloy’s plasticity and strain hardening capability [25,26].
The processing route employed for manufacturing AlCoCrFeNi HEAs also plays a decisive role in determining the resulting phases, microstructure, and consequently the mechanical properties. A diverse range of processing techniques is utilized, including casting, thermomechanical processing, severe plastic deformation, powder metallurgy (PM) and additive manufacturing (AM) [6,7,20,27,28]. Conventional casting often leads to dendritic segregation and non-equilibrium solidification defects [3,27]. Processing techniques that produce fine-grained microstructures enhance mechanical strength and superplastic forming capability [20,27,29]. On the other hand, AM provides unique control over the microstructure through rapid solidification, potentially stabilizing non-equilibrium phases and hierarchical structures [7,28]. Moreover, post-processing heat treatments can further modify the phase composition, promoting mechanisms such as dynamic recrystallization [30].
Considerable research efforts have focused on understanding how variations in composition and processing, especially the proportion of FCC and BCC phases, influence the operative deformation mechanisms [18,21,31]. FCC-rich alloys generally exhibit ductile behavior dominated by dislocation slip and twinning, whereas BCC-rich alloys demonstrate higher strength but limited ductility [31,32]. The critical balance between FCC and BCC phases is a key factor in determining the strength–ductility trade-off in AlCoCrFeNi HEAs [20]. Recent crystal plasticity simulations [18] have revealed that BCC grains develop significantly higher geometrically necessary dislocation (GND) densities than adjacent FCC grains, thereby acting as dynamic strengthening networks through phase-boundary pinning effects. Stress concentrations are particularly pronounced at grain boundaries where gradients of stress, strain, and dislocation density are observed. Moreover, the activation of slip systems in neighboring FCC and BCC regions exhibits a high degree of compatibility, indicating strong mechanical coupling across phase boundaries.
Accordingly, the non-equiatomic composition Al5Co35Cr30Fe20Ni5 (at. %) was selected to stabilize a dual-phase FCC/BCC microstructure suitable for analyzing phase-specific load partitioning and interface-mediated plasticity. Despite considerable progress in understanding dual-phase HEAs, a detailed investigation of the microstructural interactions at the FCC/BCC interfaces is still lacking. The present study aims to elucidate the plastic deformation compatibility between FCC and BCC phases in a non-equiatomic AlCoCrFeNi HEA. The material investigated in this study was a non-equiatomic dual-phase AlCoCrFeNi high-entropy alloy, processed by forging as-cast ingots at 1200 °C, followed by a short heat treatment at 1200 °C and subsequent water quenching. High-resolution techniques, including digital image correlation under in situ deformation (HRDIC) and transmission electron microscopy (TEM), are employed to characterize strain localization, interface-mediated mechanisms, and the influence of nanoprecipitates on dislocation motion. These experimental insights are used to rationalize and complement recent crystal plasticity simulations, providing a comprehensive understanding of phase-specific load partitioning and the mechanisms governing the strength–ductility balance in dual-phase HEAs.

2. Materials and Methods

The nominal composition of the HEA, given in atomic percentage, is Al5Co35Cr30Fe20Ni5. The proper amounts of pure, commercially available elements were melted in an induction furnace under a protective argon atmosphere to prevent oxidation. The purity of the different elements was as follows: electrolytic cobalt 99.95 wt.%; aluminothermic chromium 99.32 wt.% (containing as main impurities 0.23% Fe and 0.14% Al); electrolytic nickel 99.92 wt.%; iron 99.63 wt.%; and aluminum 99.95 wt.%. Melting was carried out using an induction furnace, model VCT800V from Indutherm (Indutherm Erwärmungsanlagen GmbH, Walzbachtal, Germany), with a maximum power of 20 kW. In total, 70% of the maximum power of the furnace was selected, i.e., 14 kW, for melting. The elements were cut into small pieces with sizes below 2 cm and placed in an 800 cm3 zirconia crucible. The temperature was measured using a pyrometer, which allowed for instantaneous control of the level of power supplied in the course of heating up to 1700 °C. The heating took about 15 min, and then, this temperature was maintained for 10 min to ensure good homogeneity of the melt prior to pouring it into a cylindrical copper mold. The cast ingot was cut into slices 12 mm thick, which were hot-forged at 1200 °C down to a final thickness of 6 mm (a reduction of 50%), followed by heat treatment at 1200 °C for 15 min and subsequent water quenching.
Tensile tests were performed using an MTS-870 servo-hydraulic machine (MTS Systems Corporation, Eden Prairie, MN, USA) according to the ASTM E8M standard [33] at a strain rate of 10−3 s−1. Tensile samples were cylindrical, with a gauge length of 18 mm, a diameter of 3 mm and a radius of 4 mm, while compression samples were cylindrical, measuring 10 mm in length and 5 mm in diameter (a drawing of the tensile sample is provided in Figure S1).
Microstructural characterization of the forged and thermally treated alloy before and after plastic deformation was performed using scanning (SEM) and transmission electron (TEM) microscopes equipped with energy-dispersive spectroscopy detectors (EDS). Jeol 6500F SEM (JEOL Ltd., Tokyo, Japan) was operated at 20 kV for fracture observations and electron backscatter diffraction (EBSD) measurements. In addition, the working distance for EBSD acquisition was 15 mm and the step size during acquisition varied from 0.57 μm to 0.09 μm depending on the magnification of the image. For TEM, two microscopes were used: Jeol JEM 2100 (JEOL Ltd., Tokyo, Japan) operated at 200 kV and Jeol JEM 3000F (JEOL Ltd., Tokyo, Japan) operated at 300 kV. Samples for SEM and EBSD were prepared by mechanical polishing with alumina and colloidal silica in the final steps. EBSD data processing and analysis were performed using Aztec Crystal software version 6.1 SP1 (Oxford Instruments, Abingdon, UK). Deformed specimens for TEM were prepared by electrolytic jet polishing using a reactive mixture of 10% nitric acid in methanol at −35 °C.
Synchrotron radiation diffraction (SRD) experiments during compression tests were performed on the P07 beamline of PETRA III at the Deutsches Elektronen-Synchrotron (DESY, Hamburg, Germany). In situ compressive tests were carried out at a strain rate of 10−3 s−1. Compressive samples were machined from the forged plate in the radial direction. Samples for the compression tests were cylinders with a diameter of 5 mm and a length of 10 mm. The gauge volume was defined as 0.8 × 0.8 × 5 mm3 (beam section × cylinder diameter). The diffraction patterns were obtained at a frequency of 2 Hz using a Perkin-Elmer XRD 1621 detector (PerkinElmer Inc., Waltham, MA, USA) with an array of 20482 pixels2, with an effective pixel size of 200 × 200 µm2. The wavelength was 0.0124 nm. LaB6 powders were used as a reference to calibrate the system. The detector-to-sample distance was 1646 mm. Conventional 2θ diffraction profiles were obtained by azimuthal integration of the Debye–Scherrer rings in the axial and radial directions around ±7.5°. The fitting of the diffraction peaks was carried out using a Gaussian function. The lattice strain for each orientation can be calculated by the relative shift in the position of the diffraction peak defined by
ε h k l = d h k l − d 0 , h k l d 0 , h k l
where dhkl and d0,hkl are the planar spacing of the hkl plane in the stressed and stress-free crystal. The lattice spacing and the diffraction angle θ are related through Bragg’s law throughout:
d h k l = λ 2 s i n θ h k l
Rietveld analysis is obtained from the Debye–Scherrer pattern using the MAUD software version 2.8 [34].
The characterization of the strain distribution via high-resolution digital image correlation (HRDIC) requires the development of a homogeneous and finely distributed marker pattern on the sample surface. In this study, a gold speckle pattern was obtained by remodeling a thin gold layer, previously sputtered on the sample surface, using the procedure described in [35]. Prismatic compression samples of 5 × 5 × 10 mm3 were machined with the compressive axis perpendicular to the forging direction. Then, one of the prismatic faces was ground, polished with 9, 3 and 1 µm oil-based diamond suspensions and finished using water-free fumed silica (0.2 µm) suspension. Gold speckles were generated by the flow of water vapor at 300 °C, achieving an optimum speckle pattern size of 10–30 nm after four hours of remodeling time.
Back-scattered electron images of the remodeled gold pattern were acquired using a FEI SEM FEI VERIOS 460 (FEI Company, Hillsboro, OR, USA). An area of around 91 × 61 µm2 was captured by stitching 64 images of 2048 × 1768 pixel resolution with an overlap of 20%. The samples were tested in uniaxial compression at a constant strain rate of 10−4 s−1 up to a macroscopic strain of around 2% and 4%.
Images of the same area for the undeformed and deformed states (2 and 4% plastic strain) were correlated using DaVis software version 8.3 [36], selecting a sub-window size of 8 × 8. The displacement data, with a strain resolution of 23 nm, were obtained using the same procedure described in previous papers [37].

3. Results

3.1. Microstructural Characterization

Figure 1 shows the microstructure and orientation image map (OIM) of the forged alloy in the plane formed by the forging and the radial directions. The microstructure shows a mixture of FCC and BCC phases. The microstructure is equiaxed, exhibiting similar grain sizes for both phases. The refinement of the microstructure seems to result from the recrystallization of both phases during the forging stage. Figure 1 also shows the IOM for the forged alloy for each individual phase, FCC and BCC, respectively. The volume fractions of both phases, listed in Table 1, show that the FCC phase has a higher volume fraction. Most of the FCC grains contain Σ3 annealing twin boundaries within the grains.
Figure 2a shows the Debye–Scherrer obtained from the 2D detector rings before the mechanical test. The diffraction pattern in the axial direction (Figure 2b) shows the presence of diffraction peaks corresponding to both cubic phases. Moreover, a small peak at 2θ = 2.48°, corresponding to the B2 phase is also observed. The volume fraction of each phase as well as their corresponding lattice parameters were determined through Rietveld analysis and compared with the values obtained by EBSD. As an example, the experimental and fitting curves for axial (polar angle 0°) and radial (polar angle 90°) directions in Figure 2c,d show good agreement between both curves. Table 1 lists the lattice parameters and volume fractions of the three phases. The volume fractions agree rather well with the values calculated from EBSD measurements. The only difference is the presence of the B2 phase, which is not detectable in the EBSD analyses because of their nanometric size. It is important to point out that the presence of 2% of an ordered B2 phase is, a priori, related to the BCC structure. The lattice parameters of the three phases agree with previous results reported in the literature [38].
Figure 3a shows a bright-field image of both phases, FCC and BCC, before plastic deformation. Within the FCC phase, isolated stacking faults generated during forging at high temperatures are observed (see white arrows). Moreover, in FCC grains, the presence of recrystallized Σ3 twins generated during forging is observed. Figure 3b shows a bright-field image of twin boundaries within the recrystallized FCC phase at a B=[110] zone axis for both areas around the twin boundary. The selected area electron diffraction (SAED) pattern shows the formation of double spots typical of these twins, which are also clearly seen in the simulated pattern.
On the other hand, within the BCC phase, fine spherical precipitates of around 10 nm in diameter are observed (Figure 3a). The selected area diffraction pattern (SAED) at a B=[001] zone axis reveals the presence of superlattice spots due to semicoherent B2 precipitates. Figure 4a shows an HRTEM image of the BCC grain at a B=[001] zone axis, demonstrating the semicoherency between B2 precipitates and the BCC matrix, as indicated by the superlattice spots. The composition of B2 precipitates and the BCC matrix was analyzed to study their differences (Figure 4b). The BCC matrix is basically composed of Fe, Cr and Co (50%Cr-30%Co-20%Fe) without Ni and Al. The precipitates are composed of five elements (30% Co, 15% Ni, 15% Fe, 15% Cr, and 25% Al). It is important to note that the external enrichment of iron in the precipitates is balanced by nickel depletion in the same extension, which indicates that the distribution of elements is not homogeneous in B2 precipitates.
Metals 16 00300 g003
Figure 3. (a) Bright-field TEM image of the alloy prior to plastic deformation, showing both FCC and BCC phases. Isolated stacking faults generated during high-temperature forging are observed within FCC grains (indicated by white arrows). Bright-field and dark-field images obtained from the superlattice spot of the SAED pattern of fine spherical B2 precipitates observed within the BCC phase. (b) Bright-field image of twin boundaries developed in recrystallized FCC grains observed along the B=[110] zone axis on both sides of the twin boundary. The corresponding selected area electron diffraction (SAED) pattern shows the characteristic double-spot splitting associated with Σ3 twins, which is in good agreement with the simulated diffraction pattern.
Figure 3. (a) Bright-field TEM image of the alloy prior to plastic deformation, showing both FCC and BCC phases. Isolated stacking faults generated during high-temperature forging are observed within FCC grains (indicated by white arrows). Bright-field and dark-field images obtained from the superlattice spot of the SAED pattern of fine spherical B2 precipitates observed within the BCC phase. (b) Bright-field image of twin boundaries developed in recrystallized FCC grains observed along the B=[110] zone axis on both sides of the twin boundary. The corresponding selected area electron diffraction (SAED) pattern shows the characteristic double-spot splitting associated with Σ3 twins, which is in good agreement with the simulated diffraction pattern.
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Metals 16 00300 g004
Figure 4. (a) HRTEM image of a BCC grain observed along the B=[001] zone axis, showing semicoherency between the B2 precipitates and the BCC matrix. The corresponding superlattice reflections are indicated. (b) Chemical composition analysis of the BCC matrix and B2 precipitates. The BCC matrix is mainly composed of Fe, Cr and Co, whereas the B2 precipitates contain all five elements (Co, Ni, Fe, Cr and Al). An external enrichment in Fe in the precipitates is balanced by a corresponding depletion in Ni, indicating a non-homogeneous elemental distribution within the B2 precipitates.
Figure 4. (a) HRTEM image of a BCC grain observed along the B=[001] zone axis, showing semicoherency between the B2 precipitates and the BCC matrix. The corresponding superlattice reflections are indicated. (b) Chemical composition analysis of the BCC matrix and B2 precipitates. The BCC matrix is mainly composed of Fe, Cr and Co, whereas the B2 precipitates contain all five elements (Co, Ni, Fe, Cr and Al). An external enrichment in Fe in the precipitates is balanced by a corresponding depletion in Ni, indicating a non-homogeneous elemental distribution within the B2 precipitates.
Metals 16 00300 g004

3.2. Mechanical Characterization During the Tensile Test

Figure 5a shows the evolution of true stress and work hardening (WH) as a function of true strain during the tensile test at room temperature for the forged alloy. The yield stress is around 650 MPa, and the elongation to failure is about 30%. The work-hardening exhibits two zones. After yielding, the work hardening is high (around 105 MPa), but decreases rapidly beyond 2.5% plastic strain towards an asymptotic value of 103 MPa. Observation of the fracture surface (see Figure 5b) shows a typical ductile failure consisting of many dimples, resulting from the formation, growth and coalescence of voids during the failure process. No apparent distinction between FCC and BCC regions is observed, which is indicative of good compatibility of the deformation mechanism between the two phases.
The analysis of the microstructure of broken samples reveals the microstructural changes occurring during the tensile test, as presented in Figure 6. The head of the tensile tested samples does not show preferential orientations in either phase. However, Kernel Average maps do show significant changes in the density of dislocations in the BCC and FCC phases. On one hand, the BCC phase is almost free of dislocations, with only some dislocation accumulation found in local regions of the BCC phase neighboring the FCC phase. On the other hand, the FCC phase shows large differences among grains. Some grains exhibit a large dislocation density, while other grains are almost free of them. Point-to-point misorientation was measured in both types of grains. No changes in misorientation below 0.5° are found in dislocation-free grains (see Figure 6), but grains with a high density of dislocations exhibit misorientations of up to 4° (see Figure S2). In the region deformed uniformly, an increase in the density of dislocations is noticed in all grains. Dislocations accumulate at the interface with the BCC phase as well as at grain boundaries and inside the grains. Inside the grains, a dislocation network develops, which is associated with significant changes in local misorientation. Dislocations are concentrated in regions where misorientation increases. This indicates that strain induces local changes in the orientation of the recrystallized grains, leading to the formation of subgrains whose walls act as effective obstacles for dislocation slip. Microstructural features previously described in the gauge length are also observed in the region close to the fracture, but the dislocation density is even higher.

3.3. Load Partitioning Evaluation by In Situ Diffraction Experiments During the Compression Test

In order to understand the elasto-plastic behavior of the different phases and, consequently, the load partitioning between them, the evolution of the internal strains in the three phases—FCC, BCC and B2 precipitate—was separately evaluated in situ during the compression test. Diffraction peaks from 2θ patterns were fitted individually, calculating the interplanar distance d, their intensity and the full width at half maximum (FWHM). The elastic deformation was calculated using Equation (1). Figure 7 shows the evolution of lattice strains as a function of the applied stress for the following diffraction peaks: {111}, {200}, {220} and {113} for the FCC phase; {110}, {200} and {121} for the BCC phase; and {100} for the B2 precipitates in the axial direction. The corresponding compression curves are also plotted to connect the evolution of internal strains of individual diffraction peaks with different parts of the deformation curve. Furthermore, the evolution of the intensities of the diffraction peaks studied, also in the axial direction, is displayed. It is interesting to note that the stress axis is perpendicular to the forging direction.
The aspect of the compression curves well resembles that of the tensile curve, so it could be assumed that the deformation mechanisms should be the same. The macroscopic yield stress, measured at 0.2% plastic deformation, was around 725 MPa. Below the yield stress, the internal strains for all diffraction peaks exhibit an elastic relationship between the internal strain and the applied stress. However, their slopes are different for the FCC and BCC phases, which means that both phases have anisotropic elastic behavior. Assuming iso-stress behavior, the Young modulus for each studied {hkl} family of planes can be estimated (Table 2). The iso-stress approximation assumes uniform stress in both phases, which constitutes an idealized boundary condition. In real dual-phase microstructures, local mechanical interactions and compatibility constraints may produce stress heterogeneities, so the extracted phase-specific moduli should be regarded as effective values. For the FCC phase the highest values correspond to the {111} and {311} planes, while the highest value for the BCC phase is found for the {121} planes. Consequently, these are the planes inducing the most hardening for each of these two phases, at least during the elastic regime.
The evolution of the internal strain in both phases, FCC and BCC, shows a sigmoidal shape. However, the behavior after the stress, when the curve loses its linear elastic behavior, is different not only in both phases but also for the different grain families for a given phase. These observations show marked differences in how the applied stress is redistributed among the different phases, indicating the occurrence of stress redistribution between different phases and grain families. The increase in internal strain values for the FCC phase, which is the majority phase, is too small, or even constant. This means that macroscopic yielding is controlled by this phase. The macroscopic yield stress coincides with the loss of linearity of grains oriented with their {111} planes perpendicular to the compression axis. These grains exhibit the highest integrated intensity (see Figure 7) because these FCC grains are mainly oriented with the <111> direction parallel to the compression axis. Grains oriented with {220} planes perpendicular to the compression axis lose their internal strains at lower applied stress than the other grains. These grains have a high m value (around 0.41). However, since their volume fraction is lower than that of grains with their {111} planes perpendicular to the compression direction, they do not control the onset of macroscopic plastic deformation. The other two studied orientations, i.e., grains oriented with the {200} and {311} planes perpendicular to the compression axis, show different behavior after yielding. The increase in internal strain as a function of the applied stress is higher than in the elastic regime, especially for grains oriented with {200} planes perpendicular to the compression axis. This behavior implies that these grains, called hard grains, bear an additional load coming from soft grains. At 10% strain, the internal strain is also constant. Therefore, these hard grains also deform plastically at high strength. Grains oriented with {311} planes perpendicular to the compression axis show behavior similar to that of grains oriented with {111} planes perpendicular to the compression axis. Therefore, these grains also contribute to the initial yielding of the alloy.
On the other hand, BCC grains, which are the minority phase, show a high increase in internal strain values (in absolute terms) before macroscopic yielding, coinciding with the loss of linearity of FCC-{220} grains oriented perpendicularly to the compression axis. This effect is especially pronounced in grains oriented with their {200} planes perpendicular to the compression axis. These grains double their internal strains when the applied stress increases from 500 MPa (−5000 µstrain) to 750 MPa (−10,000 µstrain). The internal stress experienced by these grains is higher than the stress that would correspond to pure elastic behavior. Therefore, it is expected that the grains of the BCC phase, especially those oriented with their {200} planes perpendicular to the compression axis, can assume part of the load from FCC grains. This effective reinforcing effect of the BCC phase, occurring even before macroscopic yielding, could explain the high work-hardening values measured during the onset of the plastic regime for strains below 2%. For higher plastic strains, the evolution of the internal strain in the BCC phase changes its convexity, tending toward an asymptotic value associated with plastic deformation in this phase as well.
It is important to point out that the evolution of the internal strain of the B2 precipitates was measured following the {100} diffraction peak because the other diffraction peaks are located at the same 2θ values as the diffraction peaks associated with the BCC phase. The evolution of the internal deformation of this peak follows the same tendency as that measured for the {200} plane of the BCC phase, diverging slightly around 750 MPa.

3.4. TEM Analysis of the Deformed Microstructure After Compression Test

Figure 8 shows the bright-field image of the FCC and BCC phases of the alloy deformed at 2% plastic strain. It is interesting to point out that both phases show noteworthy dislocation activity, especially at the interfaces. Figure 9 shows the bright-field image at the zone axis B=[220] in the FCC phase. The microstructure reveals the presence of intersecting slip bands and stacking faults associated with the movement of a/3[211] Shockley dislocations. The high density of planar defects and their interaction suggest significant plastic deformation accommodated through dislocation glide on multiple slip systems. Figure 9 also shows a detail of the interactions between the slip bands and stacking faults. The SAED pattern reveals the formation of diffuse streaks perpendicular to them (white arrow in Figure 9). Therefore, it is expected that the slip bands visible in the image are themselves stacking faults oriented perpendicular to the electron beam. These stacking faults not only reflect the activity of partial dislocations typical of low stacking fault energy FCC alloys, but also act as barriers to dislocation motion, in a manner analogous to Lomer–Cottrell locks. Their interaction with gliding dislocations impedes further slip and contributes to strain hardening, highlighting their role in controlling the plastic response of the FCC phase.
Figure 10a shows a bright-field image of the BCC phase showing the dislocation interaction with B2 nanoparticles through the Orowan mechanism. Dislocations are frequently observed to be arranged in paired configurations (Figure 10b). Dislocation loops are individually observed when the dislocations bypass B2 precipitates. These B2 particles act as effective obstacles to dislocation motion, promoting Orowan looping and contributing to strengthening through precipitation hardening. The local accumulation of dislocations near the B2 interfaces highlights their role in impeding plastic flow and enhancing the mechanical stability of the BCC phase within the dual-phase alloy. These dislocation loops induce additional stress in the B2 precipitates, resulting in an increase in the internal strains of the {100} diffraction peaks (see Figure 7), which is separated from the {200} diffraction peak of the BCC, especially above 750 MPa, when the BCC yields. Moreover, the FWHM also increases with an increase in the macroscopic HEA strain, as shown in Figure 11. This implies that dislocation Orowan loops are forming around the precipitates.

3.5. Strain Compatibility Throughout HRDIC by In Situ Diffraction Experiments During the Compression Test

Figure 12 shows EBSD and HRDIC maps from the selected region of interest (ROI) after 2 and 4% macroscopic plastic strain. The strain partitioning was assessed by means of HRDIC, as it provides a discretized measurement of the in-plane displacement field at different macroscopic strain steps. It shows effective shear strain (γeff) values for two successive macroscopic strain steps, at 2 and 4%, respectively. Furthermore, Figure 12 presents the strain increases between these two strain steps as the difference between their respective γeff values. To relate the observed strain values to the corresponding FCC or BCC phases, OIM images for each phase in the studied area are represented, respectively. The γeff maps show that the strain is located essentially along macro bands aligned at ±45° with respect to the loading direction (image perpendicular axis), which are formed by individual slip bands, following a very heterogeneous distribution in both phases.
At an initial 2% plastic strain, most of the deformation is concentrated along the coarse FCC grains (highlight in light blue in the top region and green in the bottom left of the figure). Their spatial distribution within the material, with many of them in contact with each other, makes strain propagation easier and promotes the formation of the observed macro bands. At this strain level, only a few slip bands are present in some BCC grains, primarily nucleated to relieve stress at grain boundaries shared with FCC grains. After a total plastic deformation of 4%, strain intensifies along the FCC grains and becomes particularly pronounced in the BCC grains due to the nucleation of numerous slip bands. This is further evidenced in Figure 12, where the observed strain values reflect the difference between 2% and 4% deformation. Regarding strain partitioning, the results indicate greater strain accumulation in the FCC phase compared to the BCC phase. This is supported by the strain histograms presented in Figure 13, which show that the FCC phase exhibits a broader strain distribution and higher γeff values for both macroscopic deformation steps.
Figure 14 shows a detailed slip trace analysis of an area including both FCC and BCC grains after 4% of plastic strain. The slip systems activated in the FCC and BCC grains correspond to the <110>{111} and <111>{110}, respectively. In general, two slip systems are activated in FCC grains, particularly in regions near other FCC grains. Additionally, some boundaries show an easy FCC-FCC slip transfer (determined by the alignment of slip bands on both sides of the grain boundary), such as boundary G1(FCC)-G2(FCC) or G3(FCC)-G4(FCC). Such slip transfer is favored by the plastic compatibility between adjacent FCC grains. However, two distinct situations are commonly observed near FCC–BCC boundaries: (i) slip blocking, or (ii) the activation of a third slip system in the FCC grain that enables slip transfer from FCC to BCC. The first case occurs between the G2 (BCC) and G4 (FCC) grains, where the deformation of the FCC slip bands is blocked at the boundary, leading to an increase in local stress concentration. The second case is observed between the G6 (BCC) and G1 (FCC) grains, where a third slip system is activated in G1 (FCC), allowing slip transfer to G6 (BCC). This slip transfer mechanism helps relieve stress at the boundary, despite the plastic incompatibility between the two phases, and enables the BCC phase to accommodate part of the deformation. The contribution of the BCC phase to the accommodation of the imposed deformation occurs later than that of the FCC phase, as indicated by the origin of the BCC slip bands, which are located at the grain boundary with another FCC grain. This delayed response appears to be induced by slip transfer events from the FCC phase to the BCC phase.

4. Discussion

The two-phase structure of this alloy is similar to that of eutectic alloys with a composition close to that of the AlCoCrFeNi2.1 alloy [13,17,20,39,40,41,42]. However, the nature and volume fraction of the phases present in the Al5Co35Cr30Fe20Ni5 HEA are completely different due to the large differences in Al and Ni contents, which are much higher in eutectic alloys. Large Ni and Al contents stabilize the ordered BCC and FCC phases, promoting the stability of B2 and/or L12, respectively. These phases can appear, depending on thermal treatments, as main or secondary phases [20,43,44]. In the case of the Al5Co35Cr30Fe20Ni5 alloy, the low Ni and Al contents preclude the formation of the ordered L12 phase. Thus, the alloy is constituted by a disordered FCC phase as the main phase, comprising about 75% in volume fraction, while the disordered BCC phase (A2 structure) constitutes the secondary phase (about 24% in volume fraction). In addition, the A2-BCC phase contains a very low fraction of the ordered BCC phase (B2 structure). Nevertheless, no large compositional differences have been found between the FCC and BCC phases; the FCC phase is slightly enriched in Co and Ni with respect to the nominal composition of the alloy, while the BCC phase exhibits certain enrichment in Cr and Al. The microstructure of the alloy corresponds to that which is thermodynamically stable at the forging temperature, i.e., a two-phase microstructure, which is frozen during subsequent water quenching from forging temperature. Consequently, both phases are present with the same volume fractions as those found in the forged material. It is worth noting that quenching prevents the formation of the σ-phase, which could render the alloy brittle [45].
Tensile and compressive curves present similar aspects, and the yield stress is also similar, so it could be considered that the alloy behaves in the same way, disregarding the load direction. This behavior probably arises from the not fully recrystallized microstructure existing in the forged material, although microstructural features associated with the recrystallization of FCC and BCC phases are slightly different. KAM maps of the forged alloy reveal that some grains of the FCC phase contain a high density of dislocations, while other grains are almost dislocation-free (see Figure 6). The development of such a microstructure is not related to grain orientation, because grains with the same orientation can exhibit a range from high to low dislocation density (see OIM of the non-deformed forged alloy in Figure 6 (Heads-Zone)). On the other hand, BCC grains are totally free of dislocations, and there are some dislocations in local regions at the interface with the FCC phase. The distribution maps of recrystallized grains, shown in Figure 15, indicate that only a small fraction of the FCC phase corresponds to recrystallized grains free of dislocations, with the rest of the grain constituting a deformed substructure consisting of recovered grains or, even, regions in which the deformed structure generated during the deformation is still retained. Recrystallization of the BCC phase proceeds more favorably because the fraction of recrystallized and recovered regions is almost identical, with very few residual grains in which the deformed structure is still present. In any case, misorientation in recrystallized and recovered grains in the FCC phase is very small, below 1º, but it increases to 3.5° in deformed grains (compare misorientation plots of the forged alloy presented in Figure 6 and Figure S2).
Deformation of dual-phase AlCoCrFeNi HEAs, especially those compositions close to eutectic compositions, is very complex because a large number of deformation mechanisms have been identified: the formation of stacking faults, which can act as pinning points for dislocations, leading to dislocation entanglements [13,44,46]; the hardening effect due to nanoprecipitates of the L12 or B2 phases [42,43,44]; and the piling up of dislocations at phase interfaces or the backstresses induced by such accumulation of dislocations at these interfaces [20,39,43,46,47,48]. It is well reported that deformation begins in the FCC phase because it is the softer main phase [20,42,44,47,48,49]. Synchrotron measurements clearly demonstrate that deformation is governed by deformation in the FCC phase, but there is a significant contribution from the BCC phase. There is a load transfer from FCC grains to BCC grains, which is manifested by the large increase of internal strains in the BCC phase, especially in the (200) planes. Thus, the BCC phase assumes part of the applied load, releasing the stress level of neighboring FCC grains, even when the material is deforming in the elastic regime. A similar behavior is found for the B2 nanoparticles dispersed in BCC grains, but in this case the B2 phase would assist the BCC in withstanding higher stresses transferred from the adjacent FCC grain. TEM observations prove that the dislocation line bows around B2 nanoparticles, drawing a curved segment among them. If the stress is high enough, the dislocation can overcome the particles, leaving a dislocation loop around the particles, as observed in Figure 10.
TEM and HRDIC analysis in samples compressed during the initial stage of plastic deformation, up to 4%, show slip bands and stacking faults. At low strain (2%), the deformation localizes in very few slip bands within the FCC phase, while higher strains are required to activate slip in the BCC phase. Nevertheless, the distribution of such bands is irregular in FCC grains, indicating that dislocation motion is promoted initially in those grains that are oriented favorably for slip. Similarly, grain orientation determines the activation of a secondary slip system intersecting with primary slip bands. At this strain level, only a few slip bands are present in some BCC grains, primarily nucleated to relieve stress at grain boundaries shared with FCC grains. Only with increasing the strain is a major density of bands found within the BCC phase. Activation of slip in BCC grains appears to be induced by the stress accumulation associated with dislocation pile-up at local points of the FCC/BCC interface. This slip transfer mechanism allows the release of stresses in spite of plastic incompatibility between both phases, resulting in a load transfer from FCC grains to BCC grains. TEM observations also support this argument.
The microstructure of the FCC and BCC grains continues to evolve during the course of plastic deformation. This can be clearly seen by comparing KAM maps in the tensile sample tested up to failure with the non-deformed alloy (head of the tensile sample), as shown in Figure 9. It can be observed that dislocation activity within FCC/BCC grains results in an increase in misorientation, higher in the case of the BCC phase, leading to the format Figure 16 on of low-angle subgrains. Since the accumulation of dislocations coincides with subgrain boundaries, the latter act as effective barriers for dislocation motion. It is interesting to note that dislocations tend to be concentrated close to grains not favorably oriented for the activation of the primary slip system in both phases (see Schmid factor map of Figure 15 corresponding to the same region presented in Figure 9), even after large strains, or close to the interfaces with the BCC phase.
Metals 16 00300 g016
Figure 16. Schmid factor map corresponding to the {111}<110> slip deformation system for the FCC phase. The step size during map acquisition is 0.09 μm.
Figure 16. Schmid factor map corresponding to the {111}<110> slip deformation system for the FCC phase. The step size during map acquisition is 0.09 μm.
Metals 16 00300 g016
To gain insight into the potential of the Al5Co35Cr30Fe20Ni5, its tensile properties have been compared in Table 3 with those reported for other HEAs in the AlCoCrFeNi/AlCoFeNi systems [20,40,41,50,51,52,53,54,55,56,57,58]. The data evidence that the strength of the present HEA, processed by forging at 1200 °C, is comparable to as-cast eutectic HEAs, with the difference that the volume fraction of the BCC phase is much lower than that of the hard B2 phase in eutectic alloys. It could be expected that the optimization of the thermomechanical processing could yield a high-strength fine-grained microstructure, as found in the case of thermomechanically processed eutectic HEAs (see Table 3). Furthermore, other strategies for improving the strength of the alloy could involve those used to increase the strength of superalloys: addition of small amounts of carbon to form carbides [58,59,60,61,62] or elements such as Ti to promote the formation of hard intermetallic phases. Small Ti contents have been demonstrated to be effective in hardening AlCoCrFeTNi alloys [53,54,55]. In the same way, the dispersion of fine particles of hard phases as alumina [63] or Y2O3 [64] could also improve the mechanical properties of the alloy. All this evidence suggests that there is a large margin for strengthening this alloy.

5. Conclusions

This study investigates the plastic deformation of a dual-phase (FCC+BCC) Al5Co35Cr30Fe20Ni5 HEA. The strain localization between both phases and the influence of B2 nanoprecipitates within the BCC phase on dislocation motion have been systematically examined. Based on the present results, the following conclusions are drawn.
(1) Evolution of FCC and BCC phases is different during forging. The BCC phase recrystallizes more easily than the FCC phase. The fraction of recrystallized and recovered regions is comparable in the case of the BCC phase, while most of the FCC phase appears as a deformed substructure consisting of recovered or deformed unrecrystallized grains.
(2) Yielding of the alloy is controlled by the FCC matrix, but the BCC phase assumes a significant part of the applied load during the elastic regime.
(3) Slip commences in FCC grains favorably oriented for slip in (111) planes. The subsequent accumulation of dislocations at the FCC/BCC interfaces promotes further slip within the BCC phase.
(4) As deformation proceeds, misorientation within the FCC and BCC grains increases, inducing the formation of subgrains whose boundaries are effective in hampering dislocation motion.
(5) The combination of high strength and ductility of the forged Al5Co35Cr30Fe20Ni5 HEA is due to the activation of deformation mechanisms that allow the release of the applied load, mainly in the FCC phase and, to a lesser extent, in the BCC phase.

Supplementary Materials

The following supporting information can be downloaded at: https://www.mdpi.com/article/10.3390/met16030300/s1, Figure S1: Drawing showing the geometry and dimensions of tensile samples. Figure S2: (Left) Kernel Average Misorientation Map of FCC phase on the head zone; (Right) Misorientation angle point by point inside a deformed grain of FCC phase (red arrow).

Author Contributions

Conceptualization, P.P. and G.G.; methodology, A.S., N.S., P.P., G.G. and P.A.; validation, A.S., S.P. and J.M.; formal analysis, J.M., A.O.-C., G.G. and P.P.; investigation, P.A., E.L., G.G., J.M., R.H. and P.P.; resources, A.S. and N.S.; data curation R.H., E.L., S.P., A.O.-C. and G.G.; writing—original draft preparation, G.G.; writing—review and editing, All authors.; supervision, P.P.; funding acquisition, P.P. and G.G. All authors have read and agreed to the published version of the manuscript.

Funding

We would like to acknowledge the financial support of the Spanish Ministry of Economy and Competitiveness under projects PID2019-104382RB-I00 and PID2022-143068OB-I00. The Deutsches Elektronen-Synchrotron (DESY) is acknowledged for providing beamtime at the P07 beamline of the PETRA III synchrotron facility under proposal I-20221015EC. The research leading to this result has received funding from the European Union’s Horizon 2020 research and innovation programme under grant agreement No. 730872 (CALIPSOplus). We also acknowledge the services provided by the MiNa Laboratory at IMN, as well as funding from CM (project S2018/NMT-4291 TEC2SPACE), MINECO (project CSIC13-4E-1794) and the EU (FEDER, FSE).

Data Availability Statement

The data presented in this study are available on request from the corresponding author. The data are not publicly available due to privacy.

Acknowledgments

We would like to acknowledge the technical support of the Electron Microscopy Laboratory at the National Center for Metallurgical Research (CENIM-CSIC) for access to their facilities and expert assistance.

Conflicts of Interest

The authors declare no conflicts of interest.

Abbreviations

The following abbreviations are used in this manuscript:
HEAHigh-Entropy Alloy
FCCFace-Centered Cubic
BCCBody-Centered Cubic
HCPHexagonal Close-Packed
SROShort-Range Order
PMPowder Metallurgy
AMAdditive Manufacturing
SEMScanning Electron Microscopy
TEMTransmission Electron Microscopy
SAEDSelected Area Electron Diffraction
OIMOrientation Image Map
HRDICHigh-Resolution Digital Image Correlation
ROIRegion of Interest
KAMKernel Average Misorientation
HIPHot Isostatic Pressing
SRDSynchrotron Radiation Diffraction
XRDX-ray Diffraction
FWHMFull Width at Half Maximum
GNDGeometrically Necessary Dislocation

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Figure 1. Microstructure and orientation image map of the forged alloy in the plane defined by the forging and the radial directions. The figure includes the OIMs corresponding to both the FCC and BCC phases, presented individually. The step size during map acquisition is 0.57 μm.
Figure 1. Microstructure and orientation image map of the forged alloy in the plane defined by the forging and the radial directions. The figure includes the OIMs corresponding to both the FCC and BCC phases, presented individually. The step size during map acquisition is 0.57 μm.
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Figure 2. (a) Synchrotron radiation diffraction patterns recorded on the 2D flat-panel detector before the compressive test. (b) 2θ diffraction pattern in the axial direction before the compression test. The red dotted square marks B2 peak enlarged in the inset. Rietveld fitting of the diffraction patterns as a function of 2θ before the compression test in the (c) axial and (d) radial directions. Diamonds correspond to the experimental data while the red line corresponds to the fitted curve.
Figure 2. (a) Synchrotron radiation diffraction patterns recorded on the 2D flat-panel detector before the compressive test. (b) 2θ diffraction pattern in the axial direction before the compression test. The red dotted square marks B2 peak enlarged in the inset. Rietveld fitting of the diffraction patterns as a function of 2θ before the compression test in the (c) axial and (d) radial directions. Diamonds correspond to the experimental data while the red line corresponds to the fitted curve.
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Figure 5. (a) True stress–true strain curve and corresponding work-hardening rate as a function of true strain for the forged alloy tested in tension at room temperature. (b) Fracture surface observed after tensile failure.
Figure 5. (a) True stress–true strain curve and corresponding work-hardening rate as a function of true strain for the forged alloy tested in tension at room temperature. (b) Fracture surface observed after tensile failure.
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Figure 6. Orientation image map and kernel average misorientation of FCC and BCC phases in the tensile sample after fracture in the head (no plastic deformation), test and fracture zones. Misorientation differences in FCC and BCC grains after deformation in the head (no plastic deformation), test and fracture zones. The step size during map acquisition is 0.5 μm in the head zone and 0.09 μm in the test and fracture zones.
Figure 6. Orientation image map and kernel average misorientation of FCC and BCC phases in the tensile sample after fracture in the head (no plastic deformation), test and fracture zones. Misorientation differences in FCC and BCC grains after deformation in the head (no plastic deformation), test and fracture zones. The step size during map acquisition is 0.5 μm in the head zone and 0.09 μm in the test and fracture zones.
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Figure 7. Compressive macroscopic stress–strain curve, evolution of internal strains as a function of the applied stress for different diffraction peaks of the FCC, BCC and B2 phases, and evolution of their integrated intensity as a function of the applied stress.
Figure 7. Compressive macroscopic stress–strain curve, evolution of internal strains as a function of the applied stress for different diffraction peaks of the FCC, BCC and B2 phases, and evolution of their integrated intensity as a function of the applied stress.
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Figure 8. Bright-field TEM image showing the interphase boundary between the FCC and BCC after 2% compressive strain. The SAED1 pattern in each phase is also presented (zone axes: BFCC=[110] and BBCC=[111]).
Figure 8. Bright-field TEM image showing the interphase boundary between the FCC and BCC after 2% compressive strain. The SAED1 pattern in each phase is also presented (zone axes: BFCC=[110] and BBCC=[111]).
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Figure 9. Bright-field TEM image showing the FCC after 2% compressive strain. Detail of the dislocations and stacking faults with the SAED pattern at zone axis B=[110]. Diffuse streaks produced by stacking faults are marked with a white arrow.
Figure 9. Bright-field TEM image showing the FCC after 2% compressive strain. Detail of the dislocations and stacking faults with the SAED pattern at zone axis B=[110]. Diffuse streaks produced by stacking faults are marked with a white arrow.
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Figure 10. (a) Bright-field TEM image of the BCCthephase after 2% compressive strain. (b) Detail of the interaction between dislocations and precipitates.
Figure 10. (a) Bright-field TEM image of the BCCthephase after 2% compressive strain. (b) Detail of the interaction between dislocations and precipitates.
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Figure 11. Evolution of the FWHM of the {001}-B2 diffraction peak as a function of strain. The compressive stress–strain curve is also plotted for comparison.
Figure 11. Evolution of the FWHM of the {001}-B2 diffraction peak as a function of strain. The compressive stress–strain curve is also plotted for comparison.
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Figure 12. Effective shear strain (γeff) map after 2% and 4% of macroscopic compressive strain (and the difference between both compressive strains), together with the corresponding orientation image maps (IOM) of the FCC and BCC phases, individually. It is interesting to point out that the X-axis (horizontal) and Y-axis (vertical) are conventionally aligned.
Figure 12. Effective shear strain (γeff) map after 2% and 4% of macroscopic compressive strain (and the difference between both compressive strains), together with the corresponding orientation image maps (IOM) of the FCC and BCC phases, individually. It is interesting to point out that the X-axis (horizontal) and Y-axis (vertical) are conventionally aligned.
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Figure 13. The strain distribution for FCC, BCC and FCC+BCC grains at compressive strains of 2% and 4% and the difference between both compressive strains.
Figure 13. The strain distribution for FCC, BCC and FCC+BCC grains at compressive strains of 2% and 4% and the difference between both compressive strains.
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Figure 14. Effective shear strain (γeff) map of a detailed zone with different deformation mechanisms in the sample at 4% strain. Traces of the possible deformation systems for both FCC and BCC structures, calculated for specific grain orientations, are plotted to facilitate the identification of the active mechanisms. Grains labeled in red are FCC, while those grains labeled in black correspond to the BCC phase.
Figure 14. Effective shear strain (γeff) map of a detailed zone with different deformation mechanisms in the sample at 4% strain. Traces of the possible deformation systems for both FCC and BCC structures, calculated for specific grain orientations, are plotted to facilitate the identification of the active mechanisms. Grains labeled in red are FCC, while those grains labeled in black correspond to the BCC phase.
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Figure 15. Recrystallization grain structure of (a) FCC and (b) BCC phases. The blue color corresponds to recrystallized grains, the yellow color corresponds to a deformed substructure consisting of recovered grains and the red color corresponds to grains retaining deformation generated during forging. The step size during map acquisition is 0.5 μm.
Figure 15. Recrystallization grain structure of (a) FCC and (b) BCC phases. The blue color corresponds to recrystallized grains, the yellow color corresponds to a deformed substructure consisting of recovered grains and the red color corresponds to grains retaining deformation generated during forging. The step size during map acquisition is 0.5 μm.
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Table 1. Lattice parameters and volume fractions of the FCC, BCC and B2 phases in the forged alloy obtained using EBSD analysis and Rietveld fitting.
Table 1. Lattice parameters and volume fractions of the FCC, BCC and B2 phases in the forged alloy obtained using EBSD analysis and Rietveld fitting.
PhaseVolume Fraction (%) EBSDVolume Fraction (%) MAUDLattice Parameter, a (Å)
FCC74743.59
BCC26242.86
B2-22.87
Table 2. Young modulus calculated from Figure 7 of the studied diffraction peaks from FCC, BCC and B2 phases.
Table 2. Young modulus calculated from Figure 7 of the studied diffraction peaks from FCC, BCC and B2 phases.
FCCE (GPa)BCCE (GPa)B2E (GPa)
{111}289{110}201{100}129
{200}155{200}134  
{220}231{121}280  
{311}293    
Table 3. Comparison among the tensile properties of the present Al5Co35Cr30Fe20Ni5 HEA with other AlCoCrFeNi/AlCoFeNi HEAs. Vacuum induction melting (VIM), laser powder bed fusion (LPBF), as-cast (AC), cold-rolled (CD), annealed (AN), vacuum arc melting (VAM), arc melting (AM), powder plasma arc additive manufacturing (PPA-AD), hot-rolled (HR), aging (AG).
Table 3. Comparison among the tensile properties of the present Al5Co35Cr30Fe20Ni5 HEA with other AlCoCrFeNi/AlCoFeNi HEAs. Vacuum induction melting (VIM), laser powder bed fusion (LPBF), as-cast (AC), cold-rolled (CD), annealed (AN), vacuum arc melting (VAM), arc melting (AM), powder plasma arc additive manufacturing (PPA-AD), hot-rolled (HR), aging (AG).
ReferenceProcessingElongation (%)UTS (MPa)YS (MPa)Alloy
20VIM25.0983536AlCoCrFeNi2.1
20LPBF10.915181235AlCoCrFeNi2.1
40AC16.21050520Fe20Co20Ni41Al19
40AC+CR+AN24.215201220Fe20Co20Ni41Al19
39VAM141390 Ni30Co30Cr10Fe10Al18W2
39VAM+CR+AN171460 Ni30Co30Cr10Fe10Al18W2
39VAM+CR+AN301850 Ni30Co30Cr10Fe10Al18W2
49VIM171187550AlCrFe1.5Ni2.6
50PPA-AD35375200Al0.4CoCrFeNi
50PPA-AD+CR30650380Al0.4CoCrFeNi
51VIM14.61072545AlCoCrFeNi2.1
51VIM+HR18.51300753AlCoCrFeNi2.1
51VIM+HR+Ag15.11519951AlCoCrFeNi2.1
52AC+CR+AN58570275(FeCoNiCr)95Ti1Al4
53VAM+AN+CR+AN1315601330Ni2CoCrFeTi0.24Al0.2
54VAM+AN+CR+AN45715425Al(4at.%)CoCrFeNi
55AC62526127Al 0.25CoCrFeNi
55AC+CR2.314791280Al0.25CoCrFeNi
56AM171050620AlCoCrFeNi2.1
56AM+CR618001625AlCoCrFeNi2.1
This workForging301350650Al5Co35Cr30Fe20Ni5(at.%)
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Pérez, P.; Perosanz, S.; Medina, J.; Laurín, E.; Orozco-Caballero, A.; Hernández, R.; Stark, A.; Schell, N.; Adeva, P.; Garces, G. Load Partitioning and Strain Compatibility in a Non-Equiatomic Dual-Phase AlCoCrFeNi High-Entropy Alloy Processed by Forging. Metals 2026, 16, 300. https://doi.org/10.3390/met16030300

AMA Style

Pérez P, Perosanz S, Medina J, Laurín E, Orozco-Caballero A, Hernández R, Stark A, Schell N, Adeva P, Garces G. Load Partitioning and Strain Compatibility in a Non-Equiatomic Dual-Phase AlCoCrFeNi High-Entropy Alloy Processed by Forging. Metals. 2026; 16(3):300. https://doi.org/10.3390/met16030300

Chicago/Turabian Style

Pérez, Pablo, Sergio Perosanz, Judit Medina, Edurne Laurín, Alberto Orozco-Caballero, Rebeca Hernández, Andreas Stark, Norbert Schell, Paloma Adeva, and Gerardo Garces. 2026. "Load Partitioning and Strain Compatibility in a Non-Equiatomic Dual-Phase AlCoCrFeNi High-Entropy Alloy Processed by Forging" Metals 16, no. 3: 300. https://doi.org/10.3390/met16030300

APA Style

Pérez, P., Perosanz, S., Medina, J., Laurín, E., Orozco-Caballero, A., Hernández, R., Stark, A., Schell, N., Adeva, P., & Garces, G. (2026). Load Partitioning and Strain Compatibility in a Non-Equiatomic Dual-Phase AlCoCrFeNi High-Entropy Alloy Processed by Forging. Metals, 16(3), 300. https://doi.org/10.3390/met16030300

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