# Influence of the starting powder and ball milling on microstructure and mechanical properties of Al0.3CoCrFeNi high entropy alloy, obtained by spark plasma sintering

Adrien Saviot, Pierre Sallamand, Sophie Cazottes, Yoann Danlos, Frédéric Herbst, María del Carmen Marco de Lucas, Jean-Marie Jouvard, Sophie Le Gallet

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Adrien Saviot, Pierre Sallamand, Sophie Cazottes, Yoann Danlos, Frédéric Herbst, et al. Influence of the starting powder and ball milling on microstructure and mechanical properties of Al0.3CoCrFeNi high entropy alloy, obtained by spark plasma sintering. Materials Today Communications, 2026, 50, pp.114500. 10.1016/j.mtcomm.2025.114500. hal-05447500

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# Influence of the starting powder and ball milling on microstructure and mechanical properties of Al₀.₃CoCrFeNi high entropy alloy, obtained by spark plasma sintering

Adrien Saviot<sup>a,∗</sup> <sup> ID </sup>, Pierre Sallamand<sup>a</sup>, Sophie Cazottes<sup>b</sup>, Yoann Danlos<sup>a,c</sup>, Frédéric Herbst<sup>a</sup>, María del Carmen Marco de Lucas<sup>a</sup>, Jean-Marie Jouvard<sup>a</sup>, Sophie Le Gallet<sup>a</sup>

<figure><img src="image_2.png" /></figure>

## ARTICLE INFO

Keywords:

High Entropy Alloys (HEA)
Spark Plasma Sintering (SPS)
Mechanical activation
Core-Shell microstructure
Ultrafine grains
Mechanical properties

## ABSTRACT

This work investigates the effect of ball milling on the microstructure and mechanical properties of Al₀.₃CoCr-FeNi high entropy alloy (HEA), starting from two different powders: pre-alloyed powders obtained by gas atomization and a pure elemental mixture. The powders were densified using spark plasma sintering (SPS). For the atomized powders, ball milling leads to the formation of a heterogeneous core-shell structure that remains after sintering, with grain sizes ranging from 1 to 10 µm in the cores and less than 1 µm in the shells. Slight air contamination during the milling stage promoted the formation of nanometric AlN at grain boundaries within the shells, which inhibited grain coarsening during SPS. Additionally, Al₂O₃ precipitates formed at prior particle boundaries (PPBs). Grain size after sintering is refined from 13 µm to 3.3 µm (without and with 28 h milling of atomized powder), resulting in an increase in hardness (from 182 to 280 HV₂), yield strength (from 370 to 615 MPa), and ultimate tensile strength (from 653 to 770 MPa), accompanied by a loss in ductility (from 54 % to 18 %). However, below 28 h of milling, the strengthening is not as high as would be expected from the Hall-Petch effect. PPBs have been identified as crack initiation sites, leading to lower-than-expected tensile properties. This negative effect of PPBs can be decreased by increasing powder particle size through prolonged milling, thereby reducing PPBs density. In contrast, mechanical activation of elemental powder mixture leads to the formation of an ultrafine-grained microstructure with dispersed nanometric B₂ and Al₂O₃ phases. This elaboration route allows a higher mechanical strength (1175 MPa) and hardness (512 HV₂), but significantly lower ductility (2 %). The random distribution of minor secondary phases does not lead to PPB-related failure. This study demonstrated the strong influence of starting powders on the milling-induced microstructure, phase formation, and the resulting strengthening mechanisms governing the mechanical behavior of Al₀.₃CoCrFeNi HEAs.

## 1. Introduction

Since the pioneering work of Cantor and Yeh on the CoCrFeNi Mn composition in 2004 [1,2], high entropy alloys (HEAs) have generated a lot of investigations on new compositions [3-6]. Among these, Cantor-derived compositions are the most widely studied and generally consist in replacing Mn or adding new alloying elements to improve mechanical, corrosion or oxidation properties [3,7-9]. The AlₓCoCrFeNi family is particularly well-studied, due to its tunable microstructure as a function of Al content, going from face-centered cubic (FCC) single

phase systems to the formation of body-centered cubic (BCC), ordered BCC (B2) and sigma phases with the increase of Al content [10]. Due to its ductile FCC matrix, Al₀.₃CoCrFeNi composition is widely investigated [11-13] and can be strengthened through mechanical treatments such as cold-rolling or via the precipitation of secondary phases, including intragranular ordered FCC (L1₂) or intergranular B2 precipitates [12, 14-18]. The best mechanical results for Al₀.₃CoCrFeNi alloys produced by arc melting are achieved by combining cold-rolling with subsequent precipitation of secondary phases with a tensile yield strength that can be tuned from 160 to 1800 MPa [11,12,16,19]. Alternative processes,

<footnote><sup>a</sup> Université Bourgogne Europe, CNRS, Laboratoire Interdisciplinaire Carnot de Bourgogne ICB UMR 6303, Dijon 21000, France</footnote>

<footnote><sup>b</sup> Université de Lyon, INSA de Lyon, MATEIS, UMR CNRS 5510, Villeurbanne Cedex 69621, France</footnote>

<footnote><sup>c</sup> Université Marie et Louis Pasteur, UTBM, CNRS, Laboratoire Interdisciplinaire Carnot de Bourgogne ICB UMR 6303, Belfort 90010, France</footnote>

<footnote>* Correspondence to: 9 Avenue Alain Savary, Dijon 21000, France.</footnote>

<footnote>E-mail address: adrien.saviot@insa-lyon.fr (A. Saviot).</footnote>
such as those investigated by Muskeri, Hou and Zhao [20–22], put in evidence that combining secondary phases precipitation with a heterogeneous microstructure through thermomechanical treatment after arc-melting, results in excellent strength-ductility balance and also excellent ballistic impact resistance. This strategy of combining secondary-phase precipitation with heterogeneous grain structures has been widely studied in Cantor-like HEAs and strongly enhances their mechanical properties [23–25].

Most investigations on HEAs focus on elaboration via arc-melting. However, powder metallurgy and more particularly SPS is an interesting way to control microstructure due to its rapid process time and is suitable for HEA elaboration [26]. Compared to hot isostatic pressing, SPS cycles are shorter and samples can be much smaller, which is more suitable for screening multiple elaboration parameters, powders or compositions. In addition, the holding times used in HIP are generally long and help prevent the out-of-equilibrium conditions often observed in SPS, particularly when starting from non-pre-alloyed powders. However, in the absence of a quenching module, HIP exhibits slower cooling rates, which allows more secondary phases to precipitate during cooling compared to SPS sintering of a pre-alloyed powder. Precipitation is therefore not as well controlled as it is during thermal annealing. Gorsse [17] and Xie [27] publications on SPS of gas-atomized Al<sub>0.3</sub>CoCrFeNi powders highlighted that SPS combined with subsequent annealing treatments is promising for achieving a good compromise between mechanical strength and ductility, especially through precipitating intragranular L<sub>12</sub> phase.

Many publications already investigated the elaboration by mechanical alloying or activation of pure element mixture, followed by SPS, always leading to a major FCC structure [28–32]. What the literature emphasizes is that this elaboration route leads to differences in the nature and content of secondary phases after sintering due to the differences in the elaboration conditions: powder granularity, purity, milling cycle and balls nature and SPS parameters. In addition, heterogeneous microstructure are generally observed for HEA obtained with this elaboration route [28,31,33–42]. The microstructures are typically composed mainly of ultrafine grains and large grains at PPBs. Mechanical activation or alloying of pure elemental mixture followed by SPS generally results in ultrafine-grained microstructures with high mechanical resistance, as observed for equiatomic CoCrFeNi, reaching an ultimate tensile strength up to 978 MPa [43,44]. However, the tensile behavior of Al<sub>0.3</sub>CoCrFeNi alloys produced with this elaboration route has still not been reported in the literature.

On the other hand, mechanical milling of pure or pre-alloyed powders has proven to be an effective approach for tailoring the microstructure of various metals and alloys, in order to improve mechanical properties [45]. Indeed, plastic deformation, typically more intense at the surface of the particles, allows the formation of bimodal or harmonic microstructures, combining ultrafine and coarse grains. Such a microstructure generally exhibits a very good combination of ductility and high tensile strength [46–52]. This elaboration route have already been tested for HEA, mainly on Cantor derivate compositions [53–59]. Nagarjuna putted in evidence for Cantor alloy that compared to homogeneous structure, harmonic one increases mechanical resistance due to enhanced strain hardening [55]. Li's work on Fe<sub>50</sub>Mn<sub>30</sub>Co<sub>10</sub>Cr<sub>10</sub> highlighted that severe surface plastic deformation allows to strongly enhance both yield and ultimate tensile strengths, with only a small loss in ductility (from 16.6 % to 12.4 %) as the milling time is increased [56]. However, such elaboration route has been tested on only a few high entropy alloy compositions, and not yet on Al<sub>0.3</sub>CoCrFeNi.

The objective of this work is firstly to investigate the elaboration of heterogeneous Al<sub>0.3</sub>CoCrFeNi microstructures by sintering milled prealloyed powders. Different milling times have been tested and interrupted SPS cycles were performed to study how powder surface modifications induced by milling, influence microstructure evolution at different densification stages. As we already observed in a previous work that mechanical activation of a mixture of pure elements with the

Al<sub>0.3</sub>CoCrFeNi composition followed by subsequent SPS leads to heterogeneous FCC microstructure [28], we will discuss on the influence of the powder nature on the sintered microstructure.

Finally, the mechanical properties of the sintered Al<sub>0.3</sub>CoCrFeNi samples were evaluated by hardness and tensile tests at room temperature. The effects of key microstructural parameters such as grain size, secondary phases and prior particle boundaries (PPBs) on strengthening were also investigated.

## 2. Materials and methods

Pre-alloyed Al<sub>0.3</sub>CoCrFeNi powder was produced by gas atomization under an Ar flow using a De Laval system. The composition of the as-atomized powder is reported in Table 1. Elemental analyses were performed by inductively coupled plasma-emission spectrometry for Al, Co, Cr, Fe and Ni, by combustion-infrared absorption for C and S and by reducing melting thermal conductivity and infrared absorption for N and O respectively. Particles smaller than 45 µm were sieved and selected for the study.

The SPS and milling parameters were selected based on previous investigations on Al<sub>x</sub>CoCrFeNi at the ICB laboratory [28,60,61]. The powder was mechanically activated using a medium-energy ball mill (Fritsch Pulverisette 4). Hardened 100Cr6 steel milling bowls (250 mL) and 10 mm diameter balls were used, with a ball-to-powder mass ratio of 7:1. Milling was performed at a rotation speed of the central disk of 250 rpm and a counter-rotation of the grinding jars at -250 rpm during 4, 8, 16 and 28 h of effective milling with 30 min pauses every hour. Milled powder will be referred to as MAA-X, where X represents the milling duration (4, 8, 16 or 28 h), and A designates the as-atomized powder. Preparation of the milling vials was carried out under an argon atmosphere to limit oxygen contamination. For the route involving the mixture of pure elemental powders, the mechanical activation and the detailed analysis of the obtained powder (MAP powder) are described in [28]. Sintering was performed via SPS (FCT Systeme GmbH HPD 10) using a graphite die with an internal diameter of 30 mm under 80 MPa during the whole cycle. The sintering temperature (1100°C) was reached using a heating ramp of 50 °C/min, held for 30 s, followed by free cooling to room temperature. To limit oxygen contamination, the graphite die was filled under an argon atmosphere (O<sub>2</sub> content <1 ppm and H<sub>2</sub>O content between 5 and 9 ppm) and sintering was performed in a vacuum atmosphere (detection limit 100 Pa).

Powders and bulk HEAs surfaces were polished using SiC abrasive papers of grades P180 and 320 during 120 s, followed by polishing with 9, 3 and 1 µm diamond suspensions during 10 min at each stage. Super finishing was performed using a colloidal silica suspension of 0.03 µm. Polished surfaces were analyzed using X-Ray Diffraction (XRD) with CuKα radiation (Bruker D8 Discover). Scanning electron microscopy (SEM) was carried out at 15 kV, using a JEOL JSM 7600 F equipped with a Low Angle Backscattered Electron Detector and energy-dispersive X-ray spectroscopy (EDS) with an Oxford Instruments X-Max 80 mm² probe. Electron backscatter diffraction (EBSD) analyses were done using an EDAX-TSL Digiview camera with a 140 nm step at 20 kV. The analyses were conducted over 600 × 900 µm regions, with binning set to 13 × 13 and a scan speed of 160 fps. Data processing was conducted using the EDAX OIM 8 system. Large-area mapping resulted in substantial data points, which necessitated excluding annealing twins from grain size calculations. Consequently, twins were treated as individual grains, and the resulting grain size measurements must be considered approximate. Photoluminescence and Raman spectra were obtained with an InVia Renishaw set-up working in backscattering configuration. The wavelength was 532 nm and the excitation power was around 5 mW. Photoluminescence maps were obtained in StreamLine Mode using a x50 objective for a 48 × 33 µm² area (steps: 0.30 × 0.6 µm²) and a x20 objective for a 450 × 320 µm² area (steps: 1.0 × 3.2 µm²). Focused ion beam (FIB), using a dual beam FIB/SEM ThermoFisher Helios NanoLab600i, was used to extract a thin lamella from the sample. The
<table><caption>**Table 1**
Composition of the as-atomized Al<sub>0.3</sub>CoCrFeNi powder.</caption><thead><tr><th>Compositions (at%)</th><th>Al</th><th>Co</th><th>Cr</th><th>Fe</th><th>Ni</th><th>C</th><th>O</th><th>N</th><th>S</th></tr></thead><tbody><tr><td>Theoretical</td><td>7.00</td><td>23.25</td><td>23.25</td><td>23.25</td><td>23.25</td><td>0.0</td><td>0.0</td><td>0.0</td><td>0.0</td></tr><tr><td>Measured</td><td>7.0</td><td>23.2</td><td>23.3</td><td>23.3</td><td>23.1</td><td>0.03</td><td>0.09</td><td>0.003</td><td>0.002</td></tr></tbody></table>

microstructure characterizations were then performed by scanning transmission electron microscopy (STEM)-EDS with a JEOL JEM-2100F. Hardness measurements were performed using a Vickers ZH HD (Zwick Roell) indenter. Tensile specimens, measuring 29 mm in length with a 4 mm reduced section of 1 mm × 1 mm, were prepared by electro-erosion cutting. Only one tensile test was conducted for each sample using a Deben 2000E tensile stage at a deformation rate of 0.4 mm/s. Strain measurements were obtained from images captured every 5 s using a Hirox RH-2000 microscope.

### 3. Results and discussion

#### 3.1. Microstructure evolution during mechanical activation of pre-alloyed powder

The as-atomized powder is composed of spherical particles with equiaxed grains of around 3.5 µm (Fig. 1 - a). Minor atomization defects, such as satellites and surface layers, are observed around the particles. In the surface layers, the equiaxed grains are smaller (hundreds of nanometers). EDS mapping (Fig. 1 - a, and Table 2) reveal slight enrichment of Al and Ni at the grain boundaries (GB). The powder is predominantly FCC single-phased. A minor secondary phase is also detected (Fig. 1 - b), characterized by a single, very low-intensity peak around 44.5°, which must correspond to the Al-Ni rich B2 phase, regularly observed in Al<sub>0.3</sub>CoCrFeNi elaboration, always with only the peak (110) detected by XRD [28,29,62–64].

The powder particle size significantly increased during mechanical activation, with initial spherical particles undergoing extensive gluing, forming particles reaching sizes of up to 1 mm after 28 h of milling (Figs. 2 and 3 - a). The shape is also changing, from spherical particles to flattened ones. Ball milling also induces grain refinement, reducing grain size from few micrometers to hundreds of nanometers after 4 h. However, grain refinement is more pronounced at the particle surfaces (Fig. 2), resulting in a core-shell microstructure, which is observed even after short milling durations (Fig. 3 - a). After 28 h of milling, the cores

<table><caption>**Table 2**
EDS measurements for Al<sub>0.3</sub>CoCrFeNi atomized powder (at%).</caption><thead><tr><th></th><th>Al</th><th>Co</th><th>Cr</th><th>Fe</th><th>Ni</th></tr></thead><tbody><tr><td>Grain</td><td>6.2 ± 0.2</td><td>24.0 ± 0.3</td><td>23.0 ± 1.7</td><td>25.3 ± 1.0</td><td>21.6 ± 0.8</td></tr><tr><td>GB</td><td>10.1 ± 0.3</td><td>21.8 ± 0.8</td><td>23.3 ± 0.7</td><td>19.8 ± 0.5</td><td>25.1 ± 0.9</td></tr></tbody></table>

of the particles exhibit grains of around 30 nm against 10 nm in the boundaries. The increase in particle size due to gluing of deformed particles, which themselves exhibit core-shell microstructures, leads to the formation of a fine-grained network, as highlighted in Fig. 3 - b. Despite precautions, such as preparing the milling vials under controlled atmospheres in a glove box, slight air contamination is inevitable. Analysis of the MAA-28h powder reveals an oxygen content of 0.1 wt% compared to 0.025 wt% for the A powder.

After planetary ball milling, the FCC structure is retained as depicted in Fig. 4 - a. The XRD peaks exhibit broadening after 4 h of milling, due to severe plastic deformation caused by ball milling. This broadening continues to increase between 4 and 8 h of milling and then almost not change up to 28 h. The intensity of the (110) peak after 28 h of milling has slightly increased, which indicates an increase in secondary phase content (Fig. 4 - b).

EDS measurements of the MAA-28h powder (Table 3) reveal that the global composition is only slightly affected by the milling stage. However, the ultrafine grains in the shells exhibit iron contamination, likely originating from the 100Cr6 martensitic steel milling vials. Additionally, nanoprecipitates are observed within the particle cores, particularly after 28 h of milling. These precipitates could correspond to a B2 phase enriched in Al and Ni or to oxides regarding their darker chemical contrast (green arrows Fig. 3-k). Consequently, this increase in secondary phase content can be attributed to B2 phase precipitation and/or contamination from the hardened 100Cr6 steel milling vials.

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Morphology, microstructure, associated EDS mapping and XRD pattern of the Al<sub>0.3</sub>CoCrFeNi as-atomized powder.</figcaption></figure>
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Particles morphology (SE) and cross-sections at the center (core) and the edge (shells) of the powders (BSE) for 4 h (a-c), 8 h (d-f), 16 h (g-i) and 28 h (j-l) 28 h milled powders.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** a) Evolution of powder particle size and grain size as function of milling time and b) shell distribution in MAA-4h particle. Red arrows show shells of initial powder particles that glued together.</figcaption></figure>

## 3.2. Influence of pre-alloyed powder milling on microstructure after sintering

The A-sintered sample exhibits a single FCC phase and an equiaxed microstructure, with isotropic grain orientation (Fig. 5). Annealing twins can be observed within the grains, which account for the slight asymmetry of the (200) XRD peak [30,65,66].

In contrast, the MAA-sintered samples display core-shell microstructures. The shells consist of submicron-sized grains, while the cores exhibit grains ranging from 1 to 20 µm in size (Fig. 6). The average grain size decreases and the grain size distribution tends to narrow as the milling time of the sintered powder increases (Fig. 7). Fine grains are also observed within the cores, corresponding to the shells of the

deformed particles that have agglomerated with prolonged milling. From 16 h of mechanical activation, the cores develop a lamellar microstructure, alternating layers of grains ≤ 1 µm and layers of grains up to 10 µm (Fig. 6 - c). This lamellar texture becomes more pronounced with increasing milling time, indicating a progressive intensification of texturization during mechanical activation. Fig. 2 showed that long duration ball milling of Al<sub>0.3</sub>CoCrFeNi atomized powder tends to flatten powder particles. The observed texturing in the bulk samples probably arises from this morphological change in the particles. Similar observations were reported by Fujita et al. [67] during the grinding of pure iron powder, where plastic deformation resulted in flat powder particles displaying (001) texturing in the bulk material.

Micrometric and submicrometric precipitates are observed, mainly
<figure><img src="image_6.png" /><figcaption>Fig. 4. a) Evolution of XRD patterns for the Al<sub>0.3</sub>CoCrFeNi as-atomized with the milling time and b) superposed pattern at (111) peak for the milled powders.</figcaption></figure>

<table><caption>**Table 3**
EDS measurements on the MAA-28h powder.</caption><thead><tr><th>Composition (at %).</th><th>Al</th><th>Co</th><th>Cr</th><th>Fe</th><th>Ni</th></tr></thead><tbody><tr><td>Global</td><td>7.1<br>± 0.1</td><td>22.9<br>± 0.2</td><td>23.3<br>± 0.1</td><td>23.6<br>± 0.4</td><td>23.0<br>± 0.2</td></tr><tr><td>Core</td><td>7.2<br>± 0.1</td><td>22.6<br>± 0.4</td><td>23.3<br>± 0.2</td><td>24.1<br>± 0.9</td><td>22.9<br>± 0.3</td></tr><tr><td>Shell</td><td>6.7<br>± 0.2</td><td>19.8<br>± 0.4</td><td>22.9<br>± 0.4</td><td>30.6<br>± 1.0</td><td>20.0<br>± 0.6</td></tr></tbody></table>

at PPBs (Fig. 8) and present Al-O enrichment (Table 4), suggesting the presence of Al<sub>2</sub>O<sub>3</sub>, which is commonly formed during the sintering of Al-containing milled powders.

To confirm the nature of these precipitates within the microstructure, a polished surface was analyzed by photoluminescence and Raman spectroscopy. The photoluminescence results (Fig. 9 – a) showed a narrow emission band at specific points on the surface, consisting of two peaks at 693.0 and 694.5 nm, which are the fingerprint of Cr<sup>3+</sup> ions replacing Al<sup>3+</sup> in α-Al<sub>2</sub>O<sub>3</sub> [68]. Raman spectra (not shown here) obtained at locations with the highest red emission intensity, also display the characteristic peaks of α-Al<sub>2</sub>O<sub>3</sub> [69], in agreement with the photoluminescence results. The integrated intensity map of the emission band in Fig. 9 - b confirms that the most prominent precipitates in the microstructure are composed of α-Al<sub>2</sub>O<sub>3</sub>. Photoluminescence mapping over a large area of 450 × 320 µm<sup>2</sup> (Fig. 9 – c) shows that these precipitates are not randomly distributed, but instead tend to form near the boundaries of anterior powder particles that glued together. This is in agreement with SEM characterizations and tends to confirm that they originate from slight oxidation of powder surfaces during the ball milling stage.

Nanometric precipitates are also observed in the shells (Fig. 8). STEM-EDS characterizations were performed to identify them.

The precipitates are intergranular and intragranular. The resulting elemental mapping and EDS analyses (Fig. 10 blue arrows and Table 5) highlight that the majority of the precipitates corresponds to AlN. Their presence indicates that even if the milling jars were filled under Ar, the powder milling stage induced air contamination of the powder. This kind of contaminations are current in powder milling processes [48]. EDS also highlighted that few precipitates are Al, Ni-rich and so correspond to residual B2 precipitates. Presence of AlN at GB of the base metal and formation of Al<sub>2</sub>O<sub>3</sub> surface layer have been reported for high temperature oxidation of Al<sub>x</sub>CoCrFeNi alloys [70]. Finally minor regions presented a Cr, C enrichment that may correspond to chromium carbides coming from the 100Cr6 milling vials used. EDS mapping highlights oxygen enrichment on these regions. However as point analyses revealed few oxygen content in Cr-C rich regions, it must be due to slight surface oxidation of the carbides. The slight Fe enrichment within shells, observed before and after sintering follows this hypothesis.

Regarding the ultrafine grains in shells, several studies have previously reported such microstructures in SPSed milled powders, for example in harmonic structures where the powder only undergoes surface plastic deformation [46,51,52,71]. Other studies have linked the presence of fine grains to dynamic recrystallization at the interfaces between powder particles during SPS, a process enhanced by powder milling, due to the higher density of crystalline defects [72,73]. However, in our case, even if shells initially have smaller grains than the cores, their grain sizes fall within a similar range. This makes the significant difference in grain growth after sintering somewhat surprising. To understand the origin of this difference in grain size between cores and shells, the evolution of the microstructure during the SPS cycle was examined by means of interrupted cycles at different temperatures for the MAA-28h powder, which underwent the most extensive mechanical activation.

As depicted in Fig. 11, after sintering at 550°C, the powder retains its ultrafine grain structure, with slightly smaller grains in the shells compared to the cores. From 550°C and particularly between 660 and

<figure><img src="image_7.png" /><figcaption>Fig. 5. a) XRD pattern and b) IPF along the applied pressure direction of the SPS sample from the as-atomized powder.</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** IQ (left) and IPF (right) along the applied pressure direction for SPS-MAA- a) 4 h, b) 8 h, c) 16 h and d) 28 h samples. Green arrows highlight fine-grained regions.</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>**Fig. 7.** Grain size distribution of Al<sub>0.3</sub>CoCrFeNi SPS-A and MAA samples obtained from EBSD mappings.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Microstructure at PPBs of 1100°C SPSed MAA-28h. Yellow arrows correspond to Al-O rich precipitates.</figcaption></figure>

<table><caption>**Table 4**
EDS measurements of 1100°C SPSed MAA-28h (at%).</caption><thead><tr><th></th><th>O</th><th>Al</th><th>Co</th><th>Cr</th><th>Fe</th><th>Ni</th></tr></thead><tbody><tr><td>Global</td><td>/</td><td>9.8 ±<br>0.3</td><td>23.0 ±<br>0.3</td><td>23.4 ±<br>0.2</td><td>23.8 ±<br>0.6</td><td>23.0 ±<br>0.2</td></tr><tr><td>Core</td><td>/</td><td>6.7 ±<br>0.1</td><td>23.2 ±<br>0.1</td><td>23.5 ±<br>0.1</td><td>23.2 ±<br>0.1</td><td>23.4 ±<br>0.2</td></tr><tr><td>Shell</td><td>/</td><td>6.7 ±<br>0.1</td><td>21.9 ±<br>0.3</td><td>23.2 ±<br>0.1</td><td>26.5 ±<br>0.7</td><td>21.7 ±<br>0.4</td></tr><tr><td>Precipitates</td><td>29.1 ±<br>0.9</td><td>20.5 ±<br>0.2</td><td>12.2 ±<br>0.4</td><td>13.1 ±<br>0.3</td><td>12.9 ±<br>0.2</td><td>12.1 ±<br>0.3</td></tr></tbody></table>

870°C, in addition to AlN, which corresponds to the finest precipitates (blue arrows) observed within the shells, larger precipitates appeared mainly in the cores (red arrows). These precipitates tend to disappear above 970°C. According to XRD patterns (Fig. 11 - a), they likely correspond to the B2 phase, which is consistent with previous studies reporting the presence of B2 phase at GB in this temperature range [11, 12, 16, 18, 74, 75].

Significant grain coarsening is observed at 970°C but only inside the cores. The limited grain growth in the shells seems to come from the presence of AlN at GB, inducing Zenner pinning. After sintering at 1100°C, GB mobility increases, allowing grains to grow. This also explains the presence of some AlN precipitates inside the grains (Fig. 10). Thus, the observed grain growth heterogeneity is caused by nitrogen diffusion during milling, which led to GB pinning by AlN and not directly by the powder heterogeneous grain refinement.

### 3.3. Comparison with milling of pure elemental mixture

When sintering Al<sub>0.3</sub>CoCrFeNi from a pure elemental powder mixture, the resulting bulk microstructure (Fig. 12 - a) is highly heterogeneous with Al-Ni rich BCC regions, Co-Cr rich FCC domains and sigma phase regions due to the high Cr content. This composite microstructure is not desirable because of the significant content of sigma phase, which makes mechanical activation prior to SPS mandatory. The
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Mapping of the photoluminescence of Cr<sup>3+</sup> ions in α-Al<sub>2</sub>O<sub>3</sub> and spectra obtained at the points indicated by arrows in MAA-16h SPSed sample: a) step size x = 0.6 µm and y = 0.3 µm c) mapping with a step size of x = 1.0 µm and y = 3.2 µm.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>**Fig. 10.** a) STEM and b) EDS-mapping of MAA-28h SPSed sample. Blue, red and green arrows respectively highlight AlN, B2 precipitates and Cr carbides.</figcaption></figure>

<table><caption>**Table 5**

EDS measurements on the thin lamella of the 1100°C SPSed MAA-28h.</caption><thead><tr><th></th><th>O</th><th>N</th><th>C</th><th>Al</th><th>Co</th><th>Cr</th><th>Fe</th><th>Ni</th></tr></thead><tbody><tr><th>Grain</th><td>4.2 ± 0.6</td><td>0.0 ± 0.0</td><td>5.2 ± 2.8</td><td>5.9 ± 0.4</td><td>20.5 ± 0.8</td><td>22.1 ± 0.8</td><td>23.7 ± 1.1</td><td>18.4 ± 0.8</td></tr><tr><th>Blue</th><td>10.4 ± 3.3</td><td>20.3 ± 0.1</td><td>1.9 ± 2.2</td><td>24.7 ± 1.5</td><td>10.4 ± 1.0</td><td>11.0 ± 1.2</td><td>11.7 ± 1.1</td><td>9.5 ± 0.9</td></tr><tr><th>Red</th><td>5.6 ± 0.2</td><td>0.1 ± 0.1</td><td>3.5 ± 0.4</td><td>30.3 ± 0.6</td><td>13.0 ± 0.2</td><td>3.2 ± 0.2</td><td>10.2 ± 0.1</td><td>34.4 ± 0.9</td></tr><tr><th>Green</th><td>1.8 ± 0.1</td><td>0.0 ± 0.0</td><td>15.7 ± 2.2</td><td>1.2 ± 0.6</td><td>7.8 ± 2.8</td><td>57.0 ± 0.3</td><td>12.9 ± 0.2</td><td>12.1 ± 0.3</td></tr></tbody></table>

pure elemental mixture has also been mechanically activated and sintered, with the same milling and SPS cycles, to compare the influence of the ball milling on different starting powders. A previous investigation [28] highlighted that 28 h of planetary ball milling allowed to initiate the formation of the FCC solid solution. After SPS at 1100°C, the alloy exhibits ultrafine grains of 400 nm (Fig. 12 - b), with PPBs which present larger grains around 2.5 µm due to dynamic recrystallization and abnormal grain growth. A more detailed analysis of the microstructure evolution during SPS is detailed in [28]. Numerous nanometric precipitates at GB mainly corresponding to Al<sub>2</sub>O<sub>3</sub> and Al-Ni rich B2 (confirmed by TEM [28]) are observed. Al<sub>2</sub>O<sub>3</sub> originates from the milling stage, as also observed in the MAA samples. However, the presence of B2 phase is associated with the non-equilibrium state retained after sintering due to the rapidity of SPS. Analysis of the MAP-28h powder reveals an oxygen content of 0.61 wt% compared to 0.1 wt% for the MAA-28h powder.

Contrary to the sintered samples obtained with the MAA powder, no AlN formation and par therefore no inhibition of the grain growth has been observed at PPBs. The absence of nitrides with the reactive sintering route suggests that the Al<sub>0.3</sub>CoCrFeNi alloy is more sensitive to nitrogen diffusion than the elemental powder mixture or that the milling of the atomized powder encountered a technical problem (defective seal).

### 3.4. Influence of the microstructure on mechanical properties

#### 3.4.1. Tensile properties

The tensile curves and the evolution of mechanical properties as a function of the milling time are depicted in Fig. 13 - a and Table 6. The alloy obtained by sintering of the atomized powder exhibits a ductile behavior slightly higher than other as-SPS investigations, such as Xie [27] and Gorsse [17] but with lower yield strength and ultimate tensile strength. This difference may result from a higher average grain size in our alloy and the presence of L<sub>12</sub> precipitates for Xie and more nitrides for Gorsse. Up to 8 h of milling, the core-shell microstructures have almost no impact on ultimate tensile strength, while ductility decreases significantly. However, the yield strength increases (+22 %). Between 16 and 28 h of powder mechanical activation, the final deformation of SPS samples is almost the same. However, significant increases in both ultimate tensile strength (+14 %) and yield strength (+23 %) occur, leading to the best compromise between high mechanical resistance and ductility for the MAA samples. These values obtained are particularly noteworthy, especially for yield strength, which exceeds the results obtained by post-SPS annealing for L<sub>12</sub> precipitation (585 MPa) [17], which is currently used to achieve high resistance while maintaining a ductile behavior.

When representing the evolution of yield strength as a function of the
<figure><img src="image_13.png" /><figcaption>Fig. 11. XRD patterns of the MAA-28h powder sintered from 550 to 970°C and microstructures at a) 550, b) 660, c) 770, d) 870 and e) 970.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>Fig. 12. XRD pattern and microstructure of Al₀.₃CoCrFeNi obtained by SPS of pure elements powder mixture a) without and b) with ball milling (MAP-28h). 1 corresponds to Al-Ni rich, 2- to Co-Fe rich and 3 to Cr rich regions.</figcaption></figure>

average grain size (Fig. 13 - b), we observe that the strengthening induced by powder ball milling does not follow a simple linear fit (black dotted line on Fig. 13 - b) and so a Hall-Petch law. The need to analyze large regions to obtain a representative analysis of the whole sample is a primary factor contributing to the imprecision of the average grain size, which could affect the linear interpolation. However, if we consider the relationship established by Gorsse (red dotted line), the samples sintered from powders milled 4, 8 and 16 h deviate from the linear fit whereas the A and MAA-28h samples are in accordance. This suggests that the strengthening could be limited in the MAA samples by the presence of PPBs, which contain numerous Al₂O₃ particles. Since MAA-28h clearly

has larger powder particles, the number of PPBs is reduced, thereby limiting their influence in comparison with other MAA samples.

Conversely, MAP-28h sample obtained by reactive SPS presents almost no plastic deformation and exhibits a yield strength of 1175 MPa, surpassing Gorsse's relationship. The numerous nanometric Al₂O₃ and B₂ precipitates are not coherent with the matrix, suggesting that dislocations cannot shear them but must bypass them and can induce a supplementary strengthening (Orowan strengthening), which is not taken into account in the Hall-Petch type law. The Orowan-induced strengthening can be calculated using the following equation:
<figure><img src="image_15.png" /><figcaption>Fig. 13. a) Tensile curves and b) evolution of σ<sub>y</sub> (b) as a function of the average grain size. Black and red linear fits respectively corresponds to MAA samples and Gorsse's work [17].</figcaption></figure>

<table><caption>**Table 6**

Comparison of yield strength (σ<sub>y</sub>), ultimate tensile strength (σ<sub>UTS</sub>), elongation (A) and d (average grain size) for SPSed Al<sub>0.3</sub>CoCrFeNi from present work, Xie and Gorsse publications [17,27].</caption><thead><tr><th></th><th>σ<sub>y</sub> (MPa)</th><th>σ<sub>UTS</sub> (MPa)</th><th>A (%)</th><th>d (μm)</th></tr></thead><tbody><tr><td>Xie [27]</td><td>421</td><td>781</td><td>46</td><td>5</td></tr><tr><td>Gorsse [17]</td><td>450</td><td>800</td><td>37</td><td>9</td></tr><tr><td>A</td><td>370</td><td>653</td><td>54</td><td>13</td></tr><tr><td>MAA-4h</td><td>430</td><td>646</td><td>43</td><td>5</td></tr><tr><td>MAA-8h</td><td>452</td><td>659</td><td>29</td><td>4.3</td></tr><tr><td>MAA-16h</td><td>500</td><td>675</td><td>19</td><td>3.9</td></tr><tr><td>MAA-28h</td><td>615</td><td>770</td><td>18</td><td>3.3</td></tr><tr><td>MAP-28h</td><td></td><td>1175</td><td>2</td><td>0.9</td></tr></tbody></table>

$$ \Delta\sigma_{orowan} = \frac{0.4M}{\pi\sqrt{1-\nu}} \cdot \frac{G \cdot b}{L} \cdot \ln\left(\frac{2rp}{b}\right) \quad (1) $$

Where M is the Taylor factor (3.06), ν the Poisson coefficient (0.33), G the shear modulus of the FCC matrix (80 GPa), b the burger vector (0.254 nm) [21,22,43,63,76–78] and L the inter-particle distance:

$$ L = 2rp \left( \sqrt{\frac{\pi}{4f}} - 1 \right) \\ rp = \sqrt{2/3r} $$

Where f is the volume fraction of precipitates and r is their radius.

B2 is slightly detected by XRD (with just the (110) peak). As we approach the detection limits of XRD we can assume that its volume fraction is estimated at approximately 1 % [79]. Al₂O₃ fraction was calculated based on the percentage by weight of oxygen measured by chemical analysis and was estimated at 2 % by volume. The average precipitate diameter was previously determined by SEM and TEM [28], yielding an average value of approximately 110 nm for B2 and 190 nm for Al₂O₃. Adding Orowan's contribution, calculated from the above-assumptions on phase volume fractions, to Gorsse's relationship, a theoretical yield strength is calculated to be 1241 MPa. This value closely matches the experimental result of 1175 MPa (+ 5.6 %) and so confirms that both B2 and Al₂O₃ contribute to the material's strengthening (with respective contributions of + 80 and +78 MPa).

The fracture surfaces of the samples are shown in Fig. 14. Macro and micro-observations follow the tensile curves with A sample presenting a ductile behavior. Numerous submicrometric voids are observed, which tend to coalesce into larger micrometric ones, a characteristic of ductile

<figure><img src="image_16.png" /><figcaption>Fig. 14. Fracture surfaces of tensile A and MAA-specimens, red arrows: highlight Al₂O₃ particles and white arrows: failure at prior particle boundaries (PPBs).</figcaption></figure>
metals. When refining the grain size, the number of large voids decreases, and they are no longer visible for MAA-16 and 28 h samples. Precipitates of hundreds of nanometers appear and their number increases with milling time. These precipitates are located at the bottom of the voids. EDS measurements revealed an increase in Al-O near the precipitates, suggesting that they correspond to Al₂O₃ precipitates observed at PPBs. Macro-observations of the fracture surfaces reveal that the failure is interparticular, due to the significant presence of Al₂O₃ in the MAA samples. Even in the MAA-4h sample, signs of decohesion between the original powder particles can be observed. We observed that powder particle size increases with the milling time and therefore the Al₂O₃ rich PPB density decreases. Since these PPBs act as crack initiation sites, their higher density, for milling durations below 28 h h, is responsible for lower-than-expected yield strengths.

For the MAP-28h sample, the fracture surface shows almost no section reduction, which is indicative of brittle behavior. However, at the macroscale, no interparticular fracture is observed, consistent with a more homogeneous distribution of Al₂O₃, unlike in the MAA samples, where it trends to concentrate at PPBs. Similar to the MAA samples, many Al-O rich precipitates are visible at the bottom of the voids.

### 3.4.2. Hardness

Fig. 15 puts in evidence that increasing the milling time results in a significant increase in the hardness of Al₀.₃CoCrFeNi, from 193 HV₂ for the A sample to 280 HV₂ for the MAA-28h sample. This increase is partially attributed to grain refinement, which enhances strength via Hall-Petch effect. Compared to the results reported by Peyrouzet [63], the intrinsic hardness of our alloy is lower (100 HV vs 140 HV), likely due to a lower content of secondary phases. Indeed, this discrepancy can be explained by the most important presence of nitrides in Peyrouzet's powder, coming from the elaboration of the powder via nitrogen atomization (argon in our case). Nevertheless, the microstructure observed in our samples seems to favor grain boundary strengthening, as evidenced by a higher Hall-Petch factor (329 compared to 250 HV.µm¹/²). The linear Hall-Petch strengthening observed in the MAA samples, despite deviations in tensile testing, can be attributed to the Vickers hardness test, which induces highly localized plastic deformation and is less sensitive to weaknesses induced by PPBs.

The MAP-28h sample, obtained by reactive sintering, overpasses Hall-Petch prediction with a hardness of 512 HV₂, corroborating the additional strengthening contribution from the secondary phases.

## 4. Discussion

Regarding the presented results, opposite heterogeneous

<figure><img src="image_17.png" /><figcaption>Fig. 15. Evolution of hardness as a function of the average grain size. Black and red linear fits respectively correspond to MAA samples and Peyrouzet's work [63].</figcaption></figure>

microstructures (fine grains/coarse grains) have been obtained by sintering Al₀.₃CoCrFeNi milled pre-alloyed powders and pure element mixture. This difference is firstly caused by slight air contamination for pre-alloyed powder during milling that led to AlN formation at GB in particle's edges. Thus, grain growth during sintering was inhibited at PPBs due to GB pinning by AlN particles, leading to a core-shell microstructure, both the shells distribution and average grain size being dependent of the milling duration. In contrast, sintering of 28 h mechanically activated elemental powder mixture, yielded an ultrafine-grained microstructure. This microstructure is surrounded by micro-metric grains at PPBs due to dynamic recrystallization during SPS [28, 72]. This phenomenon was not observed for milled pre-alloyed powders. It can be attributed to the absence of numerous nanometric AlN precipitates or to the higher mechanical activation of pure powders (FWHM = 1.8° against 1.39° for MAA-28h) for equivalent milling cycle, leading to higher dislocation density which drives dynamic recrystallization [80,81]. Minor nanometric B2 precipitates were observed coming from the non-equilibrium state of the powder and the rapidity of the sintering process. Al₂O₃ precipitates are also present and both are homogeneously distributed. No AlN particles were observed, which could suggest a higher nitrogen sensitivity of the pre-alloyed powder compared to pure element mixture.

Fig. 16 presents the yield strength vs. elongation at fracture for various 3d-transition FCC high entropy alloys obtained by spark plasma sintering, without subsequent annealing treatment. It clearly highlights that milling pre-alloyed powders prior to SPS increases mechanical resistance while maintaining sufficient ductility, in contrast to milled elemental powders. When comparing results of the present work with Li and Zhang, which are to the best of our knowledge, the only published works that investigated tensile properties of HEA obtained by SPS of milled pre-alloyed powders [53,56], the yield strength appears relatively low despite the greater plastic deformation of the powder and the ductility is much more affected. Their studies concerned FCC HEA (Fe₅₀Mn₃₀Co₁₀Cr₁₀ and CoCrFeMnNi). In the present work, we used much shorter milling times (a maximum of 28 h against 150 and 100 h), a lower ball-to-powder ratio and counter-rotation of the milling vials and stage, rather than rotation of the planetary stage alone. In our case, it resulted in much higher mechanical activation, explaining the strong deformation observed in our powders. They also did not report any air contamination of the powder, and therefore did not present oxide-rich PPBs. This further supports the hypothesis that our lower-than-expected strengthening is due weakening caused by Al₂O₃ at the PPBs.

<figure><img src="image_18.png" /><figcaption>Fig. 16. Yield strength vs elongation at fracture for different 3d transition element based HEAs produced by SPS with different nature of powders [17,27, 43,44,53,56,82–85].</figcaption></figure>
## 5. Conclusions and outlooks

The impact of mechanical activation of Al<sub>0.3</sub>CoCrFeNi powders, pre-alloyed and pure elemental mixture, on the microstructures after Spark Plasma Sintering (SPS), as well as on their mechanical properties, has been investigated.

The sintering of milled atomized powders led to a core-shell microstructure, with shells presenting submicrometric grains, due to GB pinning by nanometric AlN precipitates. The longest milling duration, due to a significant increase in particle size, effectively reduced the PPB density. Al<sub>2</sub>O<sub>3</sub> particles at PPBs were always observed, independently of the milling time. The best mechanical behavior was obtained after 28 h of milling because hardness, σ<sub>γ</sub> and σ<sub>UTS</sub> increased by 45, 66 and 18 % respectively but with a ductility loss of 67 %, compared to sintered bulk of the pre-alloyed powder. This strengthening is lower than expected due to weakening caused by Al<sub>2</sub>O<sub>3</sub> precipitates at PPBs. The use of ball milled pure elements powder mixture, led to bimodal microstructure with ultrafine grains and micrometric grains at PPBs, containing minor B2 and Al<sub>2</sub>O<sub>3</sub> precipitates distributed uniformly. It exhibited the highest mechanical resistance, similar to that of cold-rolled state (1175 MPa and 512 HV<sub>2</sub>) with poor ductility (2 %).

These results demonstrate that ball milling is an interesting strategy for tailoring the Al<sub>0.3</sub>CoCrFeNi microstructure, although its effectiveness is highly dependent of the starting powders. Nevertheless, the brittle prior particle boundaries in milled atomized powders, currently limits the potential of this route. Strategies to eliminate these PPBs, such as applying pulsed currents during SPS or employing surface post-treatments like thermal plasma droplet refining before consolidation could offer promising solutions to obtain the full hardening and ductility potential of the core-shell microstructure [86,87]. Furthermore, combining both elaboration routes with established annealing treatments such as recrystallization or L1<sub>2</sub>/B2 precipitation will be essential for further optimizing the mechanical resistance/ductility balance.

### CRediT authorship contribution statement

**Adrien Saviot:** Writing – original draft, Visualization, Investigation, Conceptualization.
**Pierre Sallamand:** Writing – review & editing, Validation, Supervision, Conceptualization.
**Frédéric Herbst:** Investigation.
**Marco de Lucas María:** Investigation.
**Sophie Cazottes:** Writing – review & editing, Validation, Investigation.
**Yoann Danlos:** Investigation.
**Jouvard Jean-Marie:** Project administration, Funding acquisition.
**Sophie Le Gallet:** Writing – review & editing, Validation, Supervision, Conceptualization.

### Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgements

This work was supported by Conseil Régional de Bourgogne Franche Comté and the EUR EIPHI (ANR 17 EURE 0002 through the COCASSE project) and the FEDER.

### References

[1] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 375–377 (2004) 213–218, https://doi.org/10.1016/j.msea.2003.10.257.

[2] J.-W. Yeh, S.-K. Chen, S.-J. Lin, J.-Y. Gan, T.-S. Chin, T.-T. Shun, C.-H. Tsau, S.-Y. Chang, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299–303, https://doi.org/10.1002/adem.200300567.

[3] M. Laurent-Brocq, J.-P. Couzinié, Alliages multi-composants à haute entropie - Concepts, microstructures et propriétés mécaniques, Etude Et. propriétéS. Des. M. étaux (2018), https://doi.org/10.51257/a-v1-re269.

[4] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93, https://doi.org/10.1016/j.pmatsci.2013.10.001.

[5] Z. Lyu, C. Lee, S.-Y. Wang, X. Fan, J.-W. Yeh, P.K. Liaw, Effects of constituent elements and fabrication methods on mechanical behavior of high-entropy alloys: a review, Met. Mat. Trans. A 50 (2019) 1–28, https://doi.org/10.1007/s11661-018-4970-z.

[6] G. Talluri, R.S. Maurya, B.S. Murty, Composition design of eutectic high-entropy alloys: a review, J. Mater. Sci. 60 (2025) 1156–1186, https://doi.org/10.1007/s10853-024-10251-0.

[7] J. Liu, Z. Lv, Z. Wu, J. Zhang, C. Zheng, C. Chen, D. Ju, L. Che, Research progress on the influence of alloying elements on the corrosion resistance of high-entropy alloys, J. Alloy. Compd. 1002 (2024) 175394, https://doi.org/10.1016/j.jallcom.2024.175394.

[8] V.K. Soni, S. Sanyal, K.R. Rao, S.K. Sinha, A review on phase prediction in high entropy alloys, Proc. Inst. Mech. Eng. Part C J. Mech. Eng. Sci. 235 (2021) 6268–6286, https://doi.org/10.1177/09544062211008935.

[9] B.R. Anne, S. Shaik, M. Tanaka, A. Basu, A crucial review on recent updates of oxidation behavior in high entropy alloys, SN Appl. Sci. 3 (2021) 366, https://doi.org/10.1007/s42452-021-04374-1.

[10] M. Tokarewicz, M. Grądzka-Dahlke, Review of recent research on AlCoCrFeNi high-entropy alloy, Metals 11 (2021) 1302, https://doi.org/10.3390/met11081302.

[11] H.Y. Yasuda, H. Miyamoto, K. Cho, T. Nagase, Formation of ultrafine-grained microstructure in Al<sub>0.3</sub>CoCrFeNi high entropy alloys with grain boundary precipitates, Mater. Lett. 199 (2017) 120–123, https://doi.org/10.1016/j.matlet.2017.04.072.

[12] B. Gwalani, S. Gorse, D. Choudhuri, M. Styles, Y. Zheng, R.S. Mishra, R. Banerjee, Modifying transformation pathways in high entropy alloys or complex concentrated alloys via thermo-mechanical processing, Acta Mater. 153 (2018) 169–185, https://doi.org/10.1016/j.actamat.2018.05.009.

[13] S. Dasari, A. Sarkar, A. Sharma, B. Gwalani, D. Choudhuri, V. Soni, S. Manda, I. Samajdar, R. Banerjee, Recovery of cold-worked Al<sub>0.3</sub>CoCrFeNi complex concentrated alloy through twinning assisted B2 precipitation, Acta Mater. 202 (2021) 448–462, https://doi.org/10.1016/j.actamat.2020.10.071.

[14] M. Annasamy, N. Haghdadi, A. Taylor, P. Hodgson, D. Fabijanic, Static recrystallization and grain growth behaviour of Al<sub>0.3</sub>CoCrFeNi high entropy alloy, Mater. Sci. Eng. A 754 (2019) 282–294, https://doi.org/10.1016/j.msea.2019.03.088.

[15] S. Dasari, A. Sarkar, A. Sharma, B. Gwalani, D. Choudhuri, V. Soni, S. Manda, I. Samajdar, R. Banerjee, Recovery of cold-worked Al<sub>0.3</sub>CoCrFeNi complex concentrated alloy through twinning assisted B2 precipitation, Acta Mater. 202 (2021) 448–462, https://doi.org/10.1016/j.actamat.2020.10.071.

[16] Z. Zhu, T. Yang, R. Shi, X. Quan, J. Zhang, R. Qiu, B. Song, Q. Liu, The effects of annealing at different temperatures on microstructure and mechanical properties of cold-rolled Al<sub>0.3</sub>CoCrFeNi high-entropy alloy, Metals 11 (2021) 940, https://doi.org/10.3390/met11060940.

[17] S. Gorse, F. Peyrouzet, T. Baffie, C. Navone, J. Maisonneuve, F. Saint-Antinon, M. Descoins, K. Hoummada, R. Banerjee, A.-C. Yeh, M. Goué, Tensile properties and work hardening in Al<sub>0.3</sub>CoCrFeNi: the role of L1<sub>2</sub> precipitates and grain size, Materialia 38 (2024) 102250, https://doi.org/10.1016/j.mtla.2024.102250.

[18] E.A. Anber, A.C. Lang, E.A. Lass, P.K. Suri, J.L. Hart, D.S. D'Antuono, H. Diao, R. Feng, R. Doherty, P.K. Liaw, M.L. Taheri, Insight into the kinetic stabilization of Al<sub>0.3</sub>CoCrFeNi high-entropy alloys, Materialia 14 (2020) 100872, https://doi.org/10.1016/j.mtla.2020.100872.

[19] B. Gwalani, S. Gorse, D. Choudhuri, Y. Zheng, R.S. Mishra, R. Banerjee, Tensile yield strength of a single bulk Al<sub>0.3</sub>CoCrFeNi high entropy alloy can be tuned from 160 MPa to 1800 MPa, Scr. Mater. 162 (2019) 18–23, https://doi.org/10.1016/j scriptamat.2018.10.023.

[20] S. Muskeri, B. Gwalani, S. Jha, A. Yu, P.A. Jannotti, R.S. Haridas, B.E. Schuster, J. T. Lloyd, R.S. Mishra, S. Mukherjee, Excellent ballistic impact resistance of Al<sub>0.3</sub>CoCrFeNi multi-principal element alloy with unique bimodal microstructure, Sci. Rep. 11 (2021) 22715, https://doi.org/10.1038/s41598-021-02209-y.

[21] J. Hou, S. Wu, C. Han, Q. Fu, Y. Cai, T. Wang, C. Chen, R. Wei, F. Li, Improving the strength of Al<sub>0.3</sub>CoCrFeNi high-entropy alloy via inducing dual-heterogeneous microstructure by aging treatment, Intermetallics 173 (2024) 108403, https://doi.org/10.1016/j.intermet.2024.108403.

[22] Y. Zhao, Z. Chen, K. Yan, W. Le, S. Naseem, H. Zhang, L. Yang, Investigation on microstructure, superior tensile property and its mechanism in Al<sub>0.3</sub>CoCrFeNi high-entropy alloy via thermo-mechanical processing, Mater. Sci. Eng. A 866 (2023) 144690, https://doi.org/10.1016/j.msea.2023.144690.

[23] Y. Ma, W. Wang, J. He, Y. Zhu, X. Wu, F. Yuan, Medium/high entropy alloys with heterogeneous structures for superior properties: a review, Mater. Today Adv. 25 (2025) 100553, https://doi.org/10.1016/j.mtadv.2024.100553.

[24] P. Sathiyamoorthi, H.S. Kim, High-entropy alloys with heterogeneous microstructure: processing and mechanical properties, Prog. Mater. Sci. 123 (2022) 100709, https://doi.org/10.1016/j.pmatsci.2020.100709.

[25] D. Han, S. Cai, W. Sun, B. Yang, J. Wang, Strategies for achieving strength–ductility synergy in face-centered cubic high- and medium-entropy alloys, Adv. Eng. Mater. 25 (2023) 2300291, https://doi.org/10.1002/adem.20230291.

[26] C.O. Ujah, D.V.V. Kallon, V.S. Aigbodion, High entropy alloys prepared by spark plasma sintering: Mechanical and thermal properties, Mater. Today Sustain. 25 (2024) 100639, https://doi.org/10.1016/j.mtsust.2023.100639.
[27] S. Xie, J. Zhang, R. Li, T. Yuan, M. Zhang, M. Wang, E. Zhang, S. Deng, Q. Li, Accelerated precipitation of the B2 particles and its effect on Al0.3CoCrFeNi high-entropy alloy by electric current assisted annealing, Mater. Charact. 181 (2021) 111434, https://doi.org/10.1016/j.matchar.2021.111434.

[28] A. Saviot, P. Sallamand, J.-P. Monchoux, C. Marcelot, R. Cours, N. Geoffroy, J.-M. Jouvard, S. Le Gallet, Phase and microstructure formation during reactive spark plasma sintering of AlxCoCrFeNi (x=0.3 and 1) high entropy alloys, J. Mater. Res. Technol. 32 (2024) 3047–3059, https://doi.org/10.1016/j.jmrt.2024.08.096.

[29] R. John, A. Karati, M.M. Garlapati, M. Vaidya, R. Bhattacharya, D. Fabijanic, B. S. Murty, Influence of mechanically activated annealing on phase evolution in Al0.3CoCrFeNi high-entropy alloy, J. Mater. Sci. 54 (2019) 14588–14598, https://doi.org/10.1007/s10853-019-03917-7.

[30] B. Roy, J. Das, Strengthening face centered cubic crystals by annealing induced nano-twins, Sci. Rep. 7 (2017) 17512, https://doi.org/10.1038/s41598-017-17848-3.

[31] R. Bhattacharya, M. Annasamy, P. Cizek, M. Kamaraj, G.M. Muralikrishna, P. Hodgson, D. Fabijanic, B.S. Murty, Evolution of phase constitution with mechanical alloying and spark plasma sintering of nanocrystalline AlxCoCrFeNi (x = 0, 0.3, 0.6, 1 mol) high-entropy alloys, J. Mater. Res. 37 (2022) 959–975, https://doi.org/10.1557/s43578-021-00483-0.

[32] R. Jayasree, R.B. Mane, R. Vijay, D. Chakravarty, Effect of process control agents on mechanically alloyed Al0.3CoCrFeNi, Mater. Lett. 292 (2021) 129618, https://doi.org/10.1016/j.matlet.2021.129618.

[33] M. Tekin, G. Polat, H. Kotan, An investigation of abnormal grain growth in Zr doped CoCrFeNi HEAs through in-situ formed oxide phases, Intermetallics 146 (2022) 107588, https://doi.org/10.1016/j.internet.2022.107588.

[34] C. Wang, W. Ji, Z. Fu, Mechanical alloying and spark plasma sintering of CoCrFeNi MnAl high-entropy alloy, Adv. Powder Technol. 25 (2014) 1334–1338, https://doi.org/10.1016/j.apt.2014.03.014.

[35] S.-H. Joo, H. Kato, M.J. Jang, J. Moon, E.B. Kim, S.-J. Hong, H.S. Kim, Structure and properties of ultrafine-grained CoCrFeMnNi high-entropy alloys produced by mechanical alloying and spark plasma sintering, J. Alloy. Compd. 698 (2017) 591–604, https://doi.org/10.1016/j.jallcom.2016.12.010.

[36] W. Ji, Z. Fu, W. Wang, H. Wang, J. Zhang, Y. Wang, F. Zhang, Mechanical alloying synthesis and spark plasma sintering consolidation of CoCrFeNiAl high-entropy alloy, J. Alloy. Compd. 589 (2014) 61–66, https://doi.org/10.1016/j.jallcom.2013.11.146.

[37] H. Cheng, X. Liu, Q. Tang, W. Wang, X. Yan, P. Dai, Microstructure and mechanical properties of FeCoCrNi MnAlx high-entropy alloys prepared by mechanical alloying and hot-pressed sintering, J. Alloy. Compd. 775 (2019) 742–751, https://doi.org/10.1016/j.jallcom.2018.10.168.

[38] R. John, A. Karati, J. Joseph, D. Fabijanic, B.S. Murty, Microstructure and mechanical properties of a high entropy alloy with a eutectic composition (AlCoCrFeNi2.1) synthesized by mechanical alloying and spark plasma sintering, J. Alloy. Compd. 835 (2020) 155424, https://doi.org/10.1016/j.jallcom.2020.155424.

[39] Y.K. Yadav, M.A. Shaz, N.K. Mukhopadhyay, T.P. Yadav, High entropy alloys synthesized by mechanical alloying: a review, J. Alloy. Metall. Syst. 9 (2025) 100170, https://doi.org/10.1016/j.jalmees.2025.100170.

[40] N. Kumar, P. Sengupta, S.R. Das, S.K. Sahu, S. Pahari, A. Panigrahi, Phase and microstructure evolution in SPS-processed CoCrFeMnNi high entropy alloy: effects of heat treatment and oxidation, Met. Mater. Trans. A 56 (2025) 4391–4409, https://doi.org/10.1007/s11661-025-07917-y.

[41] S. Mohammadi, F. Akhlaghi, J.-H. Cho, Effect of sintering temperature and late PCA addition on the microstructure and mechanical behavior of a heterogeneous-multiscale AlCoCuMnNi high-entropy alloy synthesized via mechanical alloying and spark plasma sintering, Intermetallics 185 (2025) 108920, https://doi.org/10.1016/j.internet.2025.108920.

[42] P.B. Oboso, S. Oyama, J. Horioka, L.-F. Yi, T. Onda, S. Morito, Z.-C. Chen, Microstructure and mechanical properties of CuCrFeNi medium entropy alloys synthesized via mechanical alloying and spark plasma sintering, J. Alloy. Compd. 1010 (2025) 177700, https://doi.org/10.1016/j.jallcom.2024.177700.

[43] M. Li, Y. Guo, H. Wang, J. Shan, Y. Chang, Microstructures and mechanical properties of oxide dispersion strengthened CoCrFeNi high-entropy alloy produced by mechanical alloying and spark plasma sintering, Intermetallics 123 (2020) 106819, https://doi.org/10.1016/j.internet.2020.106819.

[44] I. Moravcik, J. Cizek, J. Zapletal, Z. Kovacova, J. Vesely, P. Minarik, M. Kitzmantel, E. Neubauer, I. Dlouhy, Microstructure and mechanical properties of Ni1,5Co1,5CrFeTi0,5 high entropy alloy fabricated by mechanical alloying and spark plasma sintering, Mater. Des. 119 (2017) 141–150, https://doi.org/10.1016/j.matdes.2017.01.036.

[45] K. Ameyama, F. Cazes, H. Couque, G. Dirras, S. Kikuchi, J. Li, F. Mompiou, K. Mondal, D. Orlov, B. Sharma, D. Tingaud, S.K. Vajpai, Harmonic structure, a promising microstructure design, Mater. Res. Lett. 10 (2022) 440–471, https://doi.org/10.1080/21663831.2022.2057203.

[46] G. Dirras, M. Ota, D. Tingaud, K. Ameyama, T. Sekiguchi, Microstructure evolution during direct impact loading of commercial purity α-titanium with harmonic structure design - Direct impact loading of pure harmonic α-Ti, Mat. ériaux Tech. 103 (2015) 311, https://doi.org/10.1051/mattech/2015031.

[47] L. Farbaniec, G. Dirras, A. Krawczynska, F. Mompiou, H. Couque, F. Naimi, F. Bernard, D. Tingaud, Powder metallurgy processing and deformation characteristics of bulk multimodal nickel, Mater. Charact. 94 (2014) 126–137, https://doi.org/10.1016/j.matchar.2014.05.008.

[48] P. Olier, M. Couvrat, C. Cayron, N. Lochet, L. Chaffron, Incidence of mechanical alloying contamination on oxides and carbides formation in ODS ferritic steels, J. Nucl. Mater. 442 (2013) S106–S111, https://doi.org/10.1016/j.jnucmat.2013.03.090.

[49] V. Paul, Y. Wang, K. Ameyama, M. Ota-Kawabata, T. Ohmura, Nanomechanical analysis of SUS304L stainless steel with bimodal distribution in grain size, Mater. Trans. 63 (2022) 545–554, https://doi.org/10.2320/matertrans.MT-M2021187.

[50] S.K. Vajpai, M. Ota, T. Watanabe, R. Maeda, T. Sekiguchi, T. Kusaka, K. Ameyama, The development of high performance Ti-6Al-4V alloy via a unique microstructural design with bimodal grain size distribution, Met. Mater. Trans. A 46 (2015) 903–914, https://doi.org/10.1007/s11661-014-2649-7.

[51] D. Ueda, G. Dirras, A. Hocini, D. Tingaud, M. Ameyama, P. Langlois, D. Vrel, Z. Trzaska, Data on processing of Ti-25Nb-25Zr β-titanium alloys via powder metallurgy route: methodology, microstructure and mechanical properties, Data Brief. 17 (2018) 703–708, https://doi.org/10.1016/j.dib.2018.01.093.

[52] V. Paul, Y. Wang, K. Ameyama, M. Ota-Kawabata, T. Ohmura, Nanomechanical analysis of SUS304L stainless steel with bimodal distribution in grain size, Mater. Trans. 63 (2022) 545–554, https://doi.org/10.2320/matertrans.MT-M2021187.

[53] Z. Zhang, X. Zhai, G. Chen, X. Chen, K. Ameyama, Enhanced synergy of strength-ductility and low-cycle fatigue resistance of high-entropy alloy through harmonic structure design, Scr. Mater. 213 (2022) 114591, https://doi.org/10.1016/j scriptamat.2022.114591.

[54] L. Shi, Z. Zhang, X. Chen, Fatigue crack growth behavior in CoCrFeMnNi high entropy alloy with harmonic structure topology, Int. J. Fatigue 172 (2023) 107656, https://doi.org/10.1016/j.jifatigue.2023.107656.

[55] C. Nagarjuna, J. Song, K.Y. Jeong, M.-W. Pin, G. Song, J.K. Lee, Y.-S. Na, H.S. Kim, K.B. Kim, S.-J. Hong, Developing harmonic structure in CoCrFeMnNi high entropy alloy to enhance mechanical properties via powder metallurgy approach, J. Mater. Res. Technol. 17 (2022) 1686–1695, https://doi.org/10.1016/j.jmrt.2022.01.098.

[56] G. Li, M. Liu, S. Lyu, M. Nakatani, R. Zheng, C. Ma, Q. Li, K. Ameyama, Simultaneously enhanced strength and strain hardening capacity in FeMnCoCr high-entropy alloy via harmonic structure design, Scr. Mater. 191 (2021) 196–201, https://doi.org/10.1016/j.scriptamat.2020.09.036.

[57] S. Onoue, S. Hosogi, T. Kuno, M. Kawabata, H. Fujiwara, K. Isonishi, K. Ameyama, Microstructure evolution and unique deformation behavior of a CrMnFeCoNi harmonic structure high entropy alloy at elevated temperatures, Mater. Sci. Forum 1107 (2023) 73–78, https://doi.org/10.4028/p-bkBH1.

[58] S.K. Vajpai, C. Sawangrat, O. Yamaguchi, O.P. Ciuca, K. Ameyama, Effect of bimodal harmonic structure design on the deformation behaviour and mechanical properties of Co-Cr-Mo alloy, Mater. Sci. Eng. C 58 (2016) 1008–1015, https://doi.org/10.1016/j.msec.2015.09.055.

[59] D. Banik, Neetu, S. Mukherjee, H. Fujiwara, K. Ameyama, K. Mondal, Effect of harmonic structure on the wear behavior of high entropy Cantor alloy, Wear 534–535 (2023) 205125, https://doi.org/10.1016/j.wear.2023.205125.

[60] A. Fourmont, S. Le Gallet, O. Politano, C. Desgranges, F. Baras, Effects of planetary ball milling on AlCoCrFeNi high entropy alloys prepared by Spark Plasma Sintering: experiments and molecular dynamics study, J. Alloy. Compd. 820 (2020) 153448, https://doi.org/10.1016/j.jallcom.2019.153448.

[61] A. Fourmont, S. Le Gallet, K. Houmada, M. Descoins, C. Desgranges, O. Politano, F. Baras, Effects of mechanical activation on chemical homogeneity and contamination level in dual-phase AlCoCrFeNi high entropy alloy, Mater. Chem. Phys. 272 (2021) 125000, https://doi.org/10.1016/j.matchemphys.2021.125000.

[62] S. Moemeni, K. Dehghani, Exploring the coupled effect of cold rolling level and annealing parameters on the development of microstructure, mechanical properties and texture of Al0.3CoCrFeNi high entropy alloy, Mater. Today Commun. 46 (2025) 112687, https://doi.org/10.1016/j.mtcomm.2025.112687.

[63] F. Peyrouzet, Développement de l'alliage à haute entrepie Al0.3CoCrFeNi par frittage Spark Plasma Sintering et Fusion Laser sur Lit de Poudre, These de doctorat, Bordeaux, 2022, (https://www.theses.fr/2022BORD0080) (accessed February 13, 2024).

[64] Q.H. Tang, Y. Huang, Y.Y. Huang, X.Z. Liao, T.G. Langdon, P.Q. Dai, Hardening of an Al0.3CoCrFeNi high entropy alloy via high-pressure torsion and thermal annealing, Mater. Lett. 151 (2015) 126–129, https://doi.org/10.1016/j.matlet.2015.03.066.

[65] F. Bernard, E. Gaffet, O. Boytsov, Y. Champion, 2002, Planar defects effect on XRD line profiles of Cu- nanostructured powders prepared by ball-milling.

[66] L. Balogh, G. Ribárik, T. Ungár, Stacking faults and twin boundaries in fcc crystals determined by x-ray diffraction profile analysis, J. Appl. Phys. 100 (2006) 023512, https://doi.org/10.1063/1.2216195.

[67] S. Fujita, H. Sato, S. Motozuka, Deformation type during the ball milling process: a comparative study of the microstructures formed by ball milling, uniaxial compression, and rolling, Powder Technol. 426 (2023) 118598, https://doi.org/10.1016/j.powtec.2023.118598.

[68] C. Zuo, P.W. Jagodzinski, R-line luminescence from trace amounts of Cr3+ in aluminum oxide and its dependence on sample hydration, Appl. Spectrosc. 56 (2002) 1055–1058, https://doi.org/10.1366/000370202321274872.

[69] M. Kadleiková, J. Breza, M. Veselý, Raman spectra of synthetic sapphire, Microelectron. J. 32 (2001) 955–958, https://doi.org/10.1016/S0026-2692(01)00087-8.

[70] T.M. Butler, M.J. Pavel, M.L. Weaver, The effect of annealing on the microstructures and oxidation behaviors of AlCoCrFeNi complex concentrated alloys, J. Alloy. Compd. 956 (2023) 170391, https://doi.org/10.1016/j.jallcom.2023.170391.

[71] T. Xiang, P. Du, W. Bao, Z. Cai, K. Li, G. Xie, Bimodal grain size structure design to optimize the mechanical properties of TiZrNbTa high entropy alloys/Ti composites, Mater. Sci. Eng. A 849 (2022) 143488, https://doi.org/10.1016/j.msea.2022.143488.
[72] J. Guyon, A. Hazotte, J.P. Monchoux, E. Bouzy, Effect of powder state on spark plasma sintering of TiAl alloys, Intermetallics 34 (2013) 94–100, https://doi.org/10.1016/j.intermet.2012.11.005.

[73] D. Garbiec, V. Leshchynsky, A. García-Junceda, R. Swadźba, P. Siwak, G. Adamek, K. Radwański, Microstructure and mechanical properties of spark plasma sintered and severely deformed AA7075 alloy, Metals 11 (2021) 1433, https://doi.org/10.3390/met11091433.

[74] Z.J. Shi, Z.B. Wang, X.D. Wang, S. Zhang, Y.G. Zheng, Effect of thermally induced B2 phase on the corrosion behavior of an Al0.3CoCrFeNi high entropy alloy, J. Alloy. Compd. 903 (2022) 163886, https://doi.org/10.1016/j.jallcom.2022.163886.

[75] W. Wang, W. Zhou, S. Song, K.M. Reddy, X. Wang, Effect of deformation induced B2 precipitates on the microstructure and mechanical property of Al0.3CoCrFeNi high-entropy alloy, J. Alloy. Compd. 821 (2020) 153445, https://doi.org/10.1016/j.jallcom.2019.153445.

[76] Q. Wang, Y. Ma, B. Jiang, X. Li, Y. Shi, C. Dong, P.K. Liaw, A cuboidal B2 nanoprecipitation-enhanced body-centered-cubic alloy Al0.7CoCrFe2Ni with prominent tensile properties, Scr. Mater. 120 (2016) 85–89, https://doi.org/10.1016/j.scriptamat.2016.04.014.

[77] X. Liu, H. Cheng, Z. Li, H. Wang, F. Chang, W. Wang, Q. Tang, P. Dai, Microstructure and mechanical properties of FeCoCrNi MnTi0.1C0.1 high-entropy alloy produced by mechanical alloying and vacuum hot pressing sintering, Vacuum 165 (2019) 297–304, https://doi.org/10.1016/j.vacuum.2019.04.043.

[78] J.M. Park, J. Choe, J.G. Kim, J.W. Bae, J. Moon, S. Yang, K.T. Kim, J.-H. Yu, H. S. Kim, Superior tensile properties of 1%C-CoCrFeMnNi high-entropy alloy additively manufactured by selective laser melting, Mater. Res. Lett. 8 (2020) 1–7, https://doi.org/10.1080/21663831.2019.1638844.

[79] J.-W. Yeh, S.-Y. Chang, Y.-D. Hong, S.-K. Chen, S.-J. Lin, Anomalous decrease in X-ray diffraction intensities of Cu-Ni-Al-Co-Cr-Fe-Si alloy systems with multi-principal elements, Mater. Chem. Phys. 103 (2007) 41–46, https://doi.org/10.1016/j.matchemphys.2007.01.003.

[80] J. Humphreys, G.S. Rohrer, A. Rollett, Chapter 7 - Recrystallization of Single-Phase Alloys, in: J. Humphreys, G.S. Rohrer, A. Rollett (Eds.), Recrystallization and Related Annealing Phenomena, Third Edition, Elsevier, Oxford, 2017, pp. 245–304, https://doi.org/10.1016/B978-0-08-098235-9.00007-0.

[81] K. Huang, R.E. Logé, A review of dynamic recrystallization phenomena in metallic materials, Mater. Des. 111 (2016) 548–574, https://doi.org/10.1016/j.matdes.2016.09.012.

[82] G. Li, X. Meng, C. Geng, C. Wang, H. Ren, X. Guo, S. Li, Y. Tao, Microstructure and properties of AlxCr1-xCoFeNi high-entropy alloys prepared by spark plasma sintering, Materials 18 (2025) 755, https://doi.org/10.3390/ma18040755.

[83] S. Venkatesh Kumaran, D. Garbiec, J.M. Torralba, A novel and sustainable method to develop non-equiatomic CoCrFeNi Mox high-entropy alloys via spark plasma sintering using commercial commodity powders and evaluation of its mechanical behaviour, Mater. Sci. Eng. A 878 (2023) 145207, https://doi.org/10.1016/j.msea.2023.145207.

[84] J. Zhang, L. Huang, K. Xiong, X. Wang, Z. Wang, D. Guo, Z. Li, W. Feng, Mechanical properties and water vapour corrosion behaviour of Al x CoCrFeNi high-entropy alloys, RSC Adv. 14 (2024) 24741–24748, https://doi.org/10.1039/D4RA03892D.

[85] F. Průsa, A. Šenková, V. Kučera, J. Capek, D. Vojtěch, Properties of a high-strength ultrafine-grained CoCrFeNi Mn high-entropy alloy prepared by short-term mechanical alloying and spark plasma sintering, Mater. Sci. Eng. A 734 (2018) 341–352, https://doi.org/10.1016/j.msea.2018.08.014.

[86] Y.-L. Kuo, K. Kakehi, Effect of the prior particle boundary on the microstructure and mechanical properties of hot-isostatic-pressed IN718 alloy, Mater. Trans. 58 (2017) 1042–1048, https://doi.org/10.2320/matertrans.m2017045.

[87] S. Qin, L. Yan, X. Zhang, Removing prior particle boundaries in a powder superalloy based on the interaction between pulsed electric current and chain-like structure, J. Mater. Sci. Technol. 87 (2021) 95–100, https://doi.org/10.1016/j.jmst.2021.03.003.