Revealing the Microstructure Evolution and Mechanical Properties of Al₂O₃-Reinforced FCC-CoCrFeMnNi Matrix Composites Fabricated via Gas Atomization and Spark Plasma Sintering

Abstract

In the present work, novel Al₂O₃ particles were used to reinforce heterogeneous CoCrFeMnNi high-entropy alloy (HEA) matrix composites with nano- (5.0 wt.%) and nano- + micro- (5.0 wt.% + 10.0 wt.%) specimens. Al₂O₃ particles were fabricated via gas atomization and spark plasma sintering. The microstructure evolution and properties, i.e., density, hardness, and room temperature compression, were systematically investigated. The results indicate that the concentration of the Cr element in the pure CoCrFeMnNi HEA and the HEA matrix composite can be effectively reduced by using a gas-atomized HEA powder as the matrix. The formation of an impurity phase can also be inhibited, while the distribution uniformity of matrix elements can be improved. The composites prepared via gas-atomized powders formed a network microstructure composed of continuous Al₂O₃-rich regions and isolated Al₂O₃-poor regions, exhibiting good plasticity and improved density. The relative densities of the pure HEA, nano- (5.0 wt.%), and nano- + micro- (5.0 wt.% + 10.0 wt.%) composites were 98.9%, 97%, and 94.1%, respectively. The results demonstrate a significant improvement in the relative densities compared to the values (97.2%, 95.7%, and 93.8%) of the composites prepared via mechanical alloying. In addition, compared to the compressive fracture strains of nano- (5.0 wt.%) and nano- + micro- (5.0 wt.% + 10.0 wt.%) composites based on the mechanically alloyed HEA powder, the values of the nano- (5.0 wt.%) and nano- + micro- (5.0 wt.% + 10.0 wt.%) specimens prepared via gas atomization and spark plasma sintering increased by 80% and 67%, respectively.

1. Introduction

In recent years, high-entropy alloys (HEAs) have attracted widespread attention for subverting the traditional concept of alloy system design [1,2,3,4]. HEAs are composed of five or more principal metallic elements, and each of them has an atomic percentage between 5% and 35% in equal or near-equal molar ratios. The new class of alloys provides an innovative avenue for realizing desirable properties, such as extraordinary corrosion resistance, mechanical properties, and high wear resistance due to the high-entropy effect, hysteretic diffusion effect, lattice distortion effect, and cocktail effect [5,6,7,8]. Among the HEAs, the CoCrFeMnNi alloy has great potential for engineering applications due to its excellent low-temperature mechanical properties and resistance to hydrogen embrittlement [9,10].

However, the CoCrFeMnNi alloy with a single face-centered cubic (FCC) solid-solution structure exhibits a high density (usually >8 g/cm³) and low strength at room temperature (yield strength, i.e., YS < 400 MPa). Moreover, its mechanical properties decrease significantly at high temperatures. When the temperature reaches 1073 K, the alloy softens and the tensile strength decreases to less than 200 MPa [11,12,13]. Therefore, finding ways to improve the mechanical properties of the FCC-CoCrFeMnNi alloy has become a hot topic in the study of HEAs. Inserting the ceramic particles (TiC, Al₂O₃, SiC, Y₂O₃, etc.) into the HEA to prepare the HEA matrix composite has been proven to be a practical method for improving its comprehensive properties at both room temperature and elevated temperatures [14,15,16,17,18,19]. Among the various ceramic reinforcement particles, Al₂O₃ is a good reinforcement particle due to its unique properties, such as low density, high hardness, low price, and, most importantly, its excellent chemical stability, which enables it to generally not react with the matrix alloy [1,20,21]. Therefore, several studies have been conducted on Al₂O₃ nano-particle-reinforced HEAs [7,18,21,22]. In our previous study, we designed and synthesized a novel CoCrFeMnNi matrix composite material with different contents of Al₂O₃ particles via mechanical alloying (MA) and spark plasma sintering (SPS). The results showed that the addition of Al₂O₃ particles can effectively refine the matrix grains and gradually increase the yield strength of the composites [7].

MA is one of the main methods for producing nano-structured HEA matrix composites, which can achieve a uniform dispersion of powders into a composite by breaking the ceramic clusters [23,24,25,26]. However, due to the large number of components of the HEA, when the original powder is mixed by high-energy ball milling to prepare the pre-alloyed powder, the alloying and homogenization times are generally longer and the resulting amount of the prepared powder is small. Moreover, the powder is easily polluted by the grinding medium or air during the grinding process [27,28]. At present, the commonly used methods of pulverization include the crushing method, physicochemical method, gas atomization (GA) method, and rotating electrode method. Compared to other manufacturing processes, GA is a common method for preparing spherical powder, which has a series of advantages, such as full alloying, good uniformity, high purity, and high processing efficiency [29,30]. The unique spherical properties of aerosol powders are also favorable for 3D printing. For example, Kunce et al. successfully prepared an AlCoCrFeNi HEA with a gas-atomized powder through laser engineering net-forming technology. Its hardness was as high as 534 HV_0.5, which was much higher than that of the AlCoCrFeNi alloy that was prepared via a casting method [31].

Hence, in the present study, a novel CoCrFeMnNi matrix composite with different sizes and contents of nano- and micro- + nano-Al₂O₃ particles was designed and synthesized via the GA and SPS methods. The effects of the Al₂O₃ particles and manufacturing methods on the microstructure evolution and mechanical properties of the HEA matrix were investigated.

2. Experimental Procedure

2.1. Materials and Processing

The CoCrFeMnNi HEA alloy and Al₂O₃-reinforced HEA matrix composites were synthesized via ball milling and SPS. The matrix powders were obtained using the GA method [32]. That is, the impact of high-speed airflow broke the high-alloy ingot into small droplets and then quickly condensed it at a high cooling rate. The obtained matrix powder, with a purity of 99.9%, was provided by Beijing Ryubon New Material Technology Co., Ltd. (Beijing, China). The nano- and micro-Al₂O₃ particle powders with an average particle size of 20–30 nm and 3–4 μm, respectively, were used as reinforcements. The purities of the nano- and micro-Al₂O₃ powders were 99.8% and 99.9%, respectively. The synthesis process involved the blending of the matrix powder and the reinforcement of the Al₂O₃ powder. In order to prevent excessive fine grinding of micro-Al₂O₃ particles during prolonged ball milling, a two-step ball milling method was selected during the preparation of the bimodal-size particle-reinforced HEA composites. In the process of ball milling, the matrix powder was mixed with the nano-Al₂O₃ powder (5 wt.%), and the powder blends of the obtained matrix powder and Al₂O₃ nano-powders (5 wt.%) were subjected to a QM-QX2 omnidirectional planetary ball mill (Nanjing Nanda Instrument Co., Ltd., Nanjing, China) with a ball to powder weight ratio of 10:1 using stainless steel balls. Meanwhile, 2.5 wt% of absolute ethanol was also added as a process control agent to effectively prevent cold welding on the tank bottom and wall. The milling experiments were conducted at 300 rpm for 10 h under a high-purity argon atmosphere. Then, the micro-Al₂O₃ powder (10 wt.%) was added, and ball milling (300 rpm/2 h) was continued to obtain a nano- and micro- (5.0 wt.% + 10.0 wt.%) dual-size particle-reinforced composite powder. The control sample was prepared using the same procedure; that is, the matrix powder and nano-Al₂O₃ powder (5 wt.%) were mixed after ball milling at 300 rpm for 12 h to obtain an N-5 composite powder.

Then, the prepared powders were placed into a graphite mold with an applied pressure of 15–20 MPa in a press machine (STP-40C, Shandong Liangchen Instrument Equipment Co., Ltd., Jinan, China), and SPS was carried out using LABOX-300 (Sinter Land Inc., Tokyo, Japan) at 1000 °C under 30 MPa for 10 min under a vacuum. The obtained specimen was a cylinder measuring 30 mm in diameter and 7 mm in height. The specific SPS sintering curve is shown in Figure 1. In the present study, the obtained composites with different contents and sizes of Al₂O₃ particles, i.e., the CoCrFeMnNi HEA, nano- (5.0 wt.%), and nano- and micro- (5.0 wt.% + 10.0 wt.%) specimens, were denoted as HEA, N-5, and (N-5) + (M-10), respectively.

2.2. Microstructure and Mechanical Property Characterization

The phase identification of the sintered bulk HEA alloy and HEA matrix composites was characterized using an X-ray diffractometer (XRD, DISCOVER A25, BRUKER, Karlsruhe, Germany) with Cu Kα radiation at a scanning rate of 2°/min and a scanning range of 30–100°. The morphology, grain size, and chemical composition of the specimens were analyzed using optical microscopy (OM, LJ-JX2030, Shenzhen Xinlongji instrument Equipment Co., Ltd., Shenzhen, China) and a field-emission scanning electron microscope (SEM, ZEISS Gemini 500, Carl Zeiss AG, Ostalbkreis, Germany) with energy-dispersive spectroscopy (EDS) and backscattered electron (BSE) imaging. A finer microstructure characterization of the nano-scale phases was additionally carried out with a field-emission high-resolution scanning transmission electron microscope (Cs-corrected dedicated STEM, FEI Talos F200X, FEI company, Hillsboro, OR, USA) equipped with SE/BF/DF and EDS detectors. Thin foils measuring 3 mm in diameter were mechanically polished using combined ion milling for TEM analysis.

The actual density of the bulk HEA alloy and HEA matrix composite was measured using the Archimedes drainage method. The Vickers hardness of the bulk composite was measured using a digital microhardness tester (DIIV-1000Z, Shanghai Shangcai testermachine Co., Ltd., Shanghai, China) at a load of 200 g for 15 s. At least ten indentations were carried out for each specimen, and the average value was taken as the result. Compression tests were performed on an Instron 3382 universal testing machine (Instron, Boston, MA, USA) at a strain rate of 1 × 10⁻³ s⁻¹ with Φ4 mm × 6 mm cylindrical specimens at room temperature. In order to ensure the accuracy of the experiment, three repeated tests were conducted for each sample.

3. Results and Discussion

3.1. Microstructure Characterizations

Figure 2a presents the morphology of the CoCrFeMnNi powders prepared via the GA method. It can be seen that the majority of the powders display nearly perfect spherical features, which can facilitate the powder flow and improve the efficiency of the powder delivery [18]. Moreover, a small quantity of powders have a planetary structure on their surfaces, as indicated by the dashed yellow circles. During the process of GA, smaller droplets show a faster cooling rate and shorter solidification time. The collision between small and large particles occurs during the spray process, leading to the adhesion of the small particles onto the surface of larger particles that remain in a pasty state, ultimately resulting in the formation of a satellite structure. In addition, very few powder particles possess irregular surface edges, as indicated by the dashed red circles, which is caused by the fact that the spheroidization time of the droplets is longer than the solidification time and the powder particles solidify before spheroidization is complete. The corresponding particle size distribution is exhibited in Figure 2b; the result shows that the median particle size of the CoCrFeMnNi (GA) powder is 23 μm, with the size distribution ranging from 7 μm to 41 μm. The XRD pattern of the CoCrFeMnNi powder prepared using the GA method is presented in Figure 2c, which reveals that the diffraction peaks of the powders closely match the single FCC phase typical of CoCrFeMnNi high-entropy alloy. When the 2θ values are 43.93°, 51.24°, and 75.44°, the diffraction peaks correspond to the (111), (200), and (220) lattice planes of the FCC phase, respectively. This result indicates that the CoCrFeMnNi powder prepared via the GA method formed a single FCC solid-solution structure and was completely alloyed. The above results demonstrate that the gas-atomized CoCrFeMnNi powder used in the present research exhibits high purity, an excellent degree of sphericity, and tiny particle size, which are suitable for the preparation of HEAs and HEA-based composites.

Figure 3 shows the XRD patterns of pure CoCrFeMnNi HEA and Al₂O₃/CoCrFeMnNi HEA-based composites. It can be observed from Figure 3a that FCC remains the dominant phase both in pure CoCrFeMnNi and composite materials. With the increase in Al₂O₃ content, some weak diffraction peaks of α-Al₂O₃ appear in the N-5 (GA) and (N-5) + (M-10) (GA) composites. Figure 3b displays the magnified curves of the dashed gray area in Figure 3a. It is possible that the addition of Al₂O₃ particles makes the diffraction peaks of the FCC solid-solution phase smoother. This phenomenon is due to the addition of Al₂O₃ particles, which slightly decreases the average grain size of the composites and improves the average lattice strain [33]. The pure CoCrFeMnNi HEA (GA) bulk material and gas-atomized HEA powder exhibit the same FCC diffraction peaks, indicating that the mechanical alloy (MA) and SPS process do not alter the phase structure of the alloy.

The CoCrFeMnNi powders prepared using the GA method have primary particle boundaries (PPBs), which can hinder the formation of bulk materials and affect the uniformity of the microstructure. Due to the fast consolidation rate, the SPS process is not sufficient to eliminate the PPBs. Therefore, in order to reduce the influence of PPBs and further activate the sintering process, a specific mechanical ball milling process was carried out on the gas-atomized powder before SPS sintering, which has been previously applied in the preparation of CoCrFeMnNi HEA via the GA method [34].

Figure 4 shows the SEM micrographs of the N-5 (GA) and (N-5) + (M-10) (GA) composites after ball milling for 12 h. The Al₂O₃ particles and HEA particles in the figure are indicated by pink and green arrows, respectively. Figure 2b shows that the particle size of the gas-atomized HEA powder is about 23 μm, while the nano- and micro-sized Al₂O₃-reinforced particles are 20 nm and 3–4 μm, respectively. The size difference between the matrix and the reinforced powders will affect the homogenization process during ball milling. After 12 h of ball milling, the larger-sized CoCrFeMnNi powders become elliptical, and small flake-like powder particles adhere to their surface. This phenomenon indicates that with the accumulation of milling energy in the powder, the particles are deformed to a certain extent. On the one hand, the powder particles deform severely and may break at this stage. On the other hand, the large deformation and high-energy collisions allow the local temperature to reach as high as 500–600 °C [31]. As a result, the broken-down powder particles are easily welded to each other. Through rapid breakage and welding, the thickness and shape of the powder can be changed. At the same time, the tiny Al₂O₃ particles and broken-down CoCrFeMnNi particles both gather on the surface of the large powder particles. In the N-5 (GA) and (N-5) + (M-10) (GA) composite powders, the large-sized HEA flakes have approximately equal diameters, and the agglomeration of the powder on the HEA surface becomes more pronounced with the increase in the Al₂O₃ particle content. Short-time ball milling enables the fine Al₂O₃ particles to adhere to the surface of the large HEA particles and refines the matrix powder to a certain extent.

Figure 5 shows the morphological characteristics of the sintered bulk specimens through SEM. The SEM–BSE image in Figure 5a reveals that there are fine dark gray particles dispersed within the pure CoCrFeMnNi grains (indicated by the red arrows), which are further confirmed to be Cr-rich regions (Cr-containing carbides) according to EDS analysis. The result shows that the content of Cr is more than 50 at.% and the maximum is 65.18 at.%. However, no obvious dark gray particles can be found in the matrix of the Al₂O₃-particle-reinforced composites. This phenomenon is similar to pure HEA and its composite materials prepared through mechanical alloying (MA) [7]. During high-temperature sintering, some Al₂O₃ particles decomposed and, thus, the Al element entered the matrix grains to inhibit the formation of carbides. On the one hand, the ball milling time of the powder before sintering is rather short. On the other hand, the Cr is alloyed and does not react with C. As a result, the Cr-rich areas in the HEA matrix are significantly reduced and thus present as dispersed granular structures. This result confirms that using gas-atomized powder as the matrix can effectively reduce powder pollution and inhibit the formation of impurity phases.

It can be observed from the microstructures of the three specimens mentioned above that parts of the matrix grains are approximately spherical and parts of them are flat, which is due to the influence of mechanical milling on the morphology of the matrix grains. After 12 h of high-energy ball milling, most HEA particles exhibit an equiaxed flake-like shape (Figure 4). Compared to the untreated gas-atomized powders, the grain size of the matrix in the three specimens is reduced, and the grain size of the N-5 (GA) and (N-5) + (M-10) (GA) composites is slightly finer than that of the pure CoCrFeMnNi (GA) alloy. This can be primarily attributed to three factors: first, the powder particles with a larger surface area contain more than one nucleation site in the GA process. Second, the powder mixture is repeatedly broken by cold welding during high-energy ball milling, and this severe plastic deformation results in a grain refinement effect. Third, due to the fast heating and cooling rates and short holding time, the SPS sintering process effectively inhibits grain growth during the sintering process. Furthermore, both the Al₂O₃ particles and some other oxide impurities in the composites effectively hinder grain coarsening.

In order to further analyze the distribution of the Al₂O₃ reinforcement in the composites, SEM–BSE images of the (N-5) + (M-10) (GA) composites were obtained, and a corresponding schematic diagram was developed, as shown in Figure 6a,b, respectively. It can be observed that, as a result of the combined effect of short-term mechanical ball milling and rapid SPS sintering, the composite exhibits a microstructure comprising a three-dimensional continuous region rich in Al₂O₃ particles, along with isolated islands of Al₂O₃-poor regions due to the high volume fraction of Al₂O₃ reinforcement at the grain boundaries. Notably, the continuous Al₂O₃-rich region forms a network structure at the grain boundaries. In recent years, powder metallurgy combined with in situ manufacturing technology has been employed to fabricate aluminum-based and titanium-based composites featuring this type of continuous network microstructure, which has demonstrated exceptional tensile strength and ductility [35].

To further illustrate the microstructure of the sintered samples, the elemental composition of the specimen was characterized using SEM–EDS analysis. The corresponding EDS mapping images of the monolithic CoCrFeMnNi HEA alloy, N-5, and (N-5) + (M-10) composites are shown in Figure 7, Figure 8 and Figure 9, respectively. In the pure HEA alloy, small amounts of C and O were also detected due to the addition of anhydrous ethanol as a process control agent (PCA) during milling. The partial decomposition of the PCA introduced C impurities, while the O element originated from residual air present in the glove box. Some CoCrFeMnNi particles are not completely broken during high-energy ball milling due to their high toughness. As a result, the Al₂O₃ reinforcement is not evenly distributed in the composite and is mainly distributed among the matrix particles, as confirmed by the Al-element distribution shown in Figure 8 and Figure 9. Furthermore, the five elements (Co, Cr, Fe, Mn, and Ni) in the matrix grain areas of these three materials are uniformly distributed.

Table 1. SEM–EDS point analysis results for the monolithic CoCrFeMnNi and Al₂O₃/CoCrFeMnNi composites.

Sample	Chemical Composition (at.%)
	Cr	Co	Fe	Mn	Ni
Pure HEA	19.04	20.28	20.89	19.16	20.63
N-5	20.07	18.95	19.21	19.95	21.82
(N-5) + (M-10)	18.68	19.73	20.79	19.34	21.45

presents the EDS point analysis measurements of the matrix in the pure HEA, N-5, and (N-5) + (M-10) specimens. It can be seen that the actual composition of the five elements in the matrix is close to equiatomic ratios, which is consistent with the compositional design of the HEA and composites. With the increase in the Al₂O₃ particle content, the difference in the content of each element increases slightly but is still negligible. Since the CoCrFeMnNi powder prepared via the GA method is already alloyed, the Al₂O₃ powder mixture has almost no effect on the elemental composition of the matrix elements. Moreover, the composition uniformity of the matrix is also ensured compared to the composite with mechanically alloyed HEA powder as a matrix. The above results indicate that the mixture of HEA powder and Al₂O₃ powder was prepared through the short-time ball milling and GA method. Subsequently, the resulting powder was combined with SPS to successfully fabricate HEAs and their composites with a homogeneous matrix composition.

Figure 10 further displays the elemental distribution of the grain boundary region (Al₂O₃-enriched region) in the (N-5) + (M-10) composite using the STEM mode. It can be seen that the Al₂O₃ particles are mainly distributed at the grain boundaries, with a small amount of HEA matrix and small pores; as a result, there are three kinds of shading contrast. In the HADDF image, the bright areas represent the HEA matrix, which shows the uniform distribution of five elements without significant Cr enrichment. In the dark gray Al₂O_3-particle-distribution areas, in addition to Al and O, a small amount of Mn can be observed. This phenomenon also appears in the HEA-based composite prepared using the MA method. This is due to the low Gibbs free energy of Mn, which can easily react with the residual O₂ during the milling process to form Mn oxides. In our previous research, it was also observed that MnO combined with Al₂O_3P to form MnAl₂O₄ complex compounds [7].

Figure 11 presents the in-depth TEM analysis result of the (N-5) + (M-10) composite. Equiaxed grains containing nano-twins were observed in undeformed sintered samples, as shown in Figure 11a, and the thickness of the twin layer is approximately 11 nm. The corresponding SAED pattern of the yellow dashed rectangle in Figure 11a is shown in Figure 11b, and the result demonstrates that the nano-twins have an FCC structure with a matrix crystal zone axis of [011]_M and a twin crystal zone axis of $[0\overline{1}\overline{1}]$_T. The lattice parameter is 3.51 Å, which is consistent with parameter 3.56 Å of the (Ni, Fe) phase in the standard PDF card. The existence of nano-twins can introduce additional interfaces inside the grains to hinder the dislocation movement, thereby increasing the strength of the materials [36]. The HRTEM image and the corresponding inverse Fourier transform (IFFT) pattern in the lower right corner of Figure 11c demonstrate the existence of stacking faults (indicated by white arrows) and lattice distortion (indicated by dislocation symbols). The occurrence of (111) atomic dislocation in CoCrFeMnNi HEA is prone to occur during solidification, recrystallization, or deformation processes, leading to a significant presence of stacking faults. Similar phenomena are commonly observed in FCC solid-solution-structure alloys characterized by low stacking-fault energy. In addition, a subcircular Al₂O₃ particle (indicated by blue dashed lines) can be observed in the HRTEM image, as shown in Figure 11d. Further, HRTEM images and corresponding FFT and IFFT patterns of the region at the interface between the Al₂O₃ and HEA matrix are shown in Figure 11e. The results reveal that the mismatch in the coefficients of thermal expansion (CTE) between the Al₂O₃ particle and the HEA matrix leads to the formation of high-density dislocations around the Al₂O₃ particle, as indicated by the dislocation symbols. As a result, this dislocation further enhances the strength of the alloy.

3.2. Density and Mechanical Property Analyses

To further analyze the properties of the materials, the density and corresponding mechanical properties of HEA and HEA-based composites were determined and are shown in

Table 2. Properties of the monolithic HEA and Al₂O₃-reinforced CoCrFeMnNi HEA composites by the GA and SPS methods.

Specimen	Density (g/cm3)		Relative Density (%)	Hardness (HV0.2)	Compressive Yield Strength(MPa)	Ultimate Compressive Strength (MPa)	Fracture Strain (%)
	Theoretical	Actual
Pure HEA	8.016	7.925	98.9	232	443	1755	40
N-5	7.614	7.387	97.0	276	667	1562	36
(N-5) + (M-10)	6.920	6.510	94.1	294	706	1011	15

. Moreover, Figure 12 presents the density and relative density of the three materials produced using gas-atomized powders as raw materials. The relative densities of the pure HEA, N-5, and (N-5) + (M-10) composites are 98.9%, 97.0%, and 94.1%, respectively. This result demonstrates a significant improvement in the relative densities compared to the values (97.2%, 95.7%, and 93.8%) of the composites prepared by MA in our previous work [7].

Figure 13 shows the Vickers hardness of the materials. The values of the pure HEA, N-5, and (N-5) + (M-10) composites are 232 HV, 276 HV, and 294 HV, respectively. This result indicates that the hardness of the materials made using the GA method is significantly lower than that of the composites based on mechanically alloyed HEA powder [7]. With an increase in the content of Al₂O₃ particles, the hardness of the composite slightly increases due to the effects of grain boundary strengthening and dispersion strengthening. However, the standard deviation of the hardness values also increases. This is because the Al₂O₃ particles are mainly distributed at the grain boundaries, and the strengthening effect caused by the Orowan mechanism is slightly weak; so, the strength improvement is not obvious. That is, the distribution of Al₂O₃ particles in the material is not uniform, resulting in obvious differences in the hardness of various regions in the composite. In summary, both the grain size and the distribution of the reinforced phases are crucial factors affecting the hardness of the composites.

Previous studies have confirmed that, compared to ordinary homogeneous materials, the strength in Ti-based composites can be significantly improved by adjusting the microstructure of the network, and its high ductility can also be retained [37]. The composites prepared using gas-atomized powders as raw materials exhibit good plasticity; for example, the compression fracture strain of the N-5 and (N-5) + (M-10) composites can reach 36% and 15%, respectively. However, the compression fracture strain of the N-5 and (N-5) + (M-10) composites based on mechanically alloyed HEA powders decreases to 20% and 9%, respectively [7]. Dislocation movement and deformation-induced twins are the two main plastic deformation characteristics of the HEA matrix composites. The superior plasticity of the material is mainly attributed to three aspects: (a) the preparation advantages brought by the combination of GA, short-time ball milling, and SPS sintering, such as lower contamination risk and grain refinement; (b) the additional interface of the twin boundaries provides a stable strain-hardening source by hindering dislocation motion; and (c) the network distribution of the reinforcement can improve the plasticity of the composites. On the one hand, the retained large matrix area can passivate cracks. On the other hand, at the level of the macroscopic network structure, the state of the internal stress implies that the main crack path must expand along the network boundary, and the zigzag network boundary can reduce the rate of crack propagation.

The yield strengths of the pure HEA, N-5, and (N-5) + (M-10) composites are 443 MPa, 667 MPa, and 706 MPa, respectively. This result indicates that the yield strength of the composites increases with the increase in the Al₂O₃ particle content. The improvement in the strength can be attributed to the continuous Al₂O₃ network, which can refine the matrix grains (especially nano-particles) and the grain boundary strengthening effect caused by the twins. Micron particles mainly play the role of load transfer. In addition, the soft-phase matrix grains are surrounded by the hard-mesh Al₂O₃ particles, which generate additional elastic strain energy in the soft-phase matrix region. However, the above result shows that the second-phase particles are mainly distributed in the grain boundary, which reduces the strengthening effect caused by the Orowan mechanism.

The addition of the ceramic Al₂O₃ reinforcement can improve the hardness and compressive yield strength. However, the plasticity of the material becomes worse, and the strain-hardening ability becomes weaker. As a result, the compressive rupture strength (ultimate compress strength, i.e., UCS) is reduced. This phenomenon is different from the tensile process. The plastic material is actually difficult to crush during the compression process; so, the better the plastic material, the higher the compressive rupture strength. The tensile process is the opposite. This is also illustrated by the compressive stress–strain curves shown in Figure 13 in the revised manuscript.

Compared to the composites based on mechanically alloyed HEA powders in our previous studies [7], the strength of the materials based on gas-atomized powders is significantly lower. This is due to the fact that the HEA powders are alloyed using the GA method, and the mechanical grinding time is relatively short to avoid excessively refining the micron Al₂O₃ particles. Moreover, the homogeneous distribution of the Al₂O₃ particles in the composites based on mechanically alloyed HEA powders makes the Orowan strengthening mechanism a major contributor to the strengthening of the composites.

4. Conclusions

In this study, CoCrFeMnNi powder prepared via the GA method was selected as the matrix, and pure HEA, N-5, and (N-5) + (M-10) powders were obtained by adding different amounts of nano- and micro-Al₂O₃ powder and grinding the mixed powder at 300 rpm for 12 h. Finally, the corresponding bulk material was successfully prepared by subsequent SPS sintering. The characterizations of the relevant properties of the materials included XRD, OM, SEM, and TEM analyses of their microstructures, as well as density, hardness, and room temperature compression testing to elucidate their mechanical properties. Furthermore, comparisons were made with the materials prepared using the MA method. The main conclusions are as follows:

(1)

During the process of high-energy ball milling, the thickness and shape of the powder can be changed by the rapid crushing and welding of the powder. After 12 h of ball milling, the larger-sized HEA powder is elliptical, and the surface is attached to fine Al₂O₃ particles and the broken-down HEA powder.

(2)

The dispersion of granular Cr-rich carbides occurs in the pure HEA alloy and is not observed in the composite materials. Compared to the material based on mechanically alloyed HEA powder, the Cr-rich region of the alloy prepared using the GA method is obviously reduced. The actual composition of the five elements in the HEA matrix region is close to the equiatomic ratio, which is consistent with the compositional design. It was proven that using gas-atomized HEA powder as the matrix can effectively reduce powder pollution, inhibit the formation of the impurity phase, and improve composition uniformity.

(3)

The composites prepared using gas-atomized powders form a network microstructure composed of continuous Al₂O₃-rich regions and isolated Al₂O₃-poor regions, exhibiting improved hardness and yield strength. However, with the addition of hard ceramic Al₂O₃ reinforcement, the plasticity of the composites becomes worse and the strain-hardening ability becomes weaker; as a result, the compressive rupture strength (ultimate compress strength, i.e., UCS) of the composite is reduced.