Influence of Cooling Methods on Microstructure and Mechanical Properties of TiB₂@Ti/AlCoCrFeNi_2.1 Eutectic High-Entropy Alloy Matrix Composites

Abstract

The present study focused on 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 eutectic high-entropy alloy matrix composites (EHEAMCs), which were treated with furnace cooling (FC), air cooling (AC), and water cooling (WC) after being held at 1000 °C for 12 h, aiming to investigate the effect of cooling methods on their microstructure and mechanical properties. The results showed that the composites in all states consisted of FCC phase, BCC phase, TiB₂ phase, and Ti phase. The cooling methods did not change the phase types but affected the diffraction peak characteristics. With the increase in cooling rate, the diffraction peaks of FCC and BCC phases gradually separated from overlapping, and the diffraction peak of the FCC (111) crystal plane shifted to a lower angle (due to the increase in lattice constant caused by Ti element diffusion), while the diffraction peak intensity showed a downward trend. In terms of microstructure, all composites under the three cooling conditions were composed of eutectic matrix, solid solution zone, and grain boundary zone. The cooling rate had little effect on the morphology but significantly affected the element distribution. During slow cooling (FC, AC), Ti and B diffused sufficiently from the grain boundary to the matrix, resulting in higher concentrations of Ti and B in the matrix (Ti in FCC phase: 7.4 at.%, B in BCC phase: 8.1 at.% in FC state). During rapid cooling (WC), diffusion was inhibited, leading to lower concentrations in the matrix (Ti in FCC phase: 4.6 at.%, B in BCC phase: 4.3 at.%), but the element distribution was more uniform. Mechanical properties decreased with the increase in cooling rate: the FC state showed the optimal average hardness (627.0 ± 26.1 HV), yield strength (1574 MPa), fracture strength (2824 MPa), and fracture strain (24.2%); the WC state had the lowest performance (hardness: 543.2 ± 35.4 HV and yield strength: 1401 MPa) but was still better than the as-sintered state. Solid solution strengthening was the main mechanism, and slow cooling promoted element diffusion to enhance lattice distortion, achieving the synergistic improvement of strength and plasticity.

1. Introduction

High-entropy alloys (HEAs), characterized by the synergistic effect of multiple principal elements, have become a research hotspot in the field of materials due to their excellent mechanical properties (such as high strength and high toughness) and service stability [1,2,3,4]. Among them, eutectic high-entropy alloys (EHEAs) achieve a combination of high strength, high hardness, and good plasticity through their unique layered or network structures composed of eutectic and non-eutectic phases, showing great application potential in high-end equipment fields such as aerospace and marine engineering [5,6,7,8,9,10]. Studies have shown that AlCoCrFeNi_2.1, as a typical eutectic high-entropy alloy, has an as-cast structure consisting of a Co-Fe-Ni-rich face-centered cubic (FCC) phase and an Al-Ni-rich body-centered cubic (BCC) phase. Its room-temperature tensile strength can reach 1.2 GPa, and the elongation can reach 22.8%, but there are still problems such as low elastic limit to be solved urgently [8].

To further optimize the performance of eutectic high-entropy alloys, preparing EHEAMCs by adding reinforcing phases has become an effective approach. At present, strengthening methods mainly focus on particle reinforcement and process control. In terms of particle reinforcement, ceramic particles such as TiB₂, WC, and SiC are widely used due to their high hardness and chemical stability. Han et al. [11] found that the compressive strength of AlCoCrFeNi_2.1 composites reinforced with 5 vol.% TiB₂ can reach 2500 MPa, and the hardness is increased to 780 HV. Wang et al. [12] increased the yield strength of CoCrFeMnNi high-entropy alloy by 42% and reduced the friction coefficient by 22.4% by adding 5 wt.% TiB₂. Guo et al. [13] prepared SiC-doped AlCoCrFeNi_2.1 composites with a room-temperature tensile strength of 1466 MPa, which confirms the effectiveness of particle reinforcement. Ren et al. [14] reported that the introduction of TiB₂@Ti as a reinforcement into CoCrFeNi high-entropy alloy exhibits unique advantages: TiB₂, as a high-hardness ceramic phase (with a Vickers hardness of over 3000 HV), can significantly enhance the mechanical properties of the composite through dispersion strengthening. Ti can improve the interfacial compatibility between the reinforcement and the high-entropy alloy matrix, promote the solid solution of Ti elements into the matrix to generate solid solution strengthening, and meanwhile reduce interfacial defects, thus avoiding stress concentration caused by pure ceramic particles. This structure can realize the multi-mechanism effect of ‘particle strengthening—solid solution strengthening—interface synergy’ in the eutectic high-entropy alloy system and has good adaptability to the layered structure of the AlCoCrFeNi_2.1 eutectic matrix. It helps to improve strength while maintaining excellent plasticity; therefore, TiB₂@Ti is selected as the reinforcement phase in the present study. In terms of process control, technologies such as hot-press sintering, hot extrusion, and laser additive manufacturing can further improve performance by refining grains and optimizing interface bonding. Peng et al. [15] prepared AlCoCrFeNi_2.1 alloy by powder hot extrusion-annealing process, with a yield strength of 1.2 GPa and a uniform elongation of 18%. Liu et al. [16] regulated the interface of WC/AlCoCrFeNi_2.1 composites by rapid hot-press sintering, resulting in a hardness of 491.9 HV and a strength of 972.9 MPa.

As a key means to control material properties, heat treatment can significantly optimize the mechanical properties of high-entropy alloys and their composites by regulating atomic diffusion, phase evolution, and microstructure. Thermomechanical treatment (such as the synergy of annealing and deformation) has been proven to refine the nanostructure of AlCoCrFeNi_2.1, increasing the yield strength to 1182 MPa [8,15], while post-sintering heat treatment (such as heat preservation and cooling) affects performance by controlling phase stability and element distribution. Xi et al. [17] found that after annealing, the interface bonding strength of WC/AlCoCrFeNi_2.1 composites prepared by selective laser melting is improved, and the fracture mechanism changes from brittle fracture to ductile fracture. Guo et al. [18] studied the influence of heat treatment temperature on the microstructure and properties of TiB₂@Ti/AlCoCrFeNi_2.1 eutectic EHEAMCs and found that the heat treatment temperature did not change the phase structure of the composites. However, the increase in heat treatment temperature caused Ti elements with large atomic radii to dissolve in the FCC phase of the matrix, and at the same time, the alloying elements in the matrix diffused to the grain boundaries, thereby causing solid solution strengthening and grain boundary strengthening and synergistically improving the strength and plasticity of the composites. Among the key parameters of heat treatment, the cooling method directly affects the phase structure and tissue uniformity of materials by regulating atomic diffusion rate and crystal defect evolution and is the “last mile” regulation means for performance optimization.

The influence of cooling methods on high-entropy alloys and their composites has received preliminary attention. Zhou et al. [19] found that after CuCrFeMnNi high-entropy alloys were treated with furnace cooling (FC), air cooling (AC), and water cooling (WC), the offset of the diffraction peak position of the FCC phase increased with the increase in cooling rate, and the microhardness showed a trend of “slow cooling being higher than fast cooling”. Yao et al. [20] showed in their study on CoCrFeMnNi-Mo₅C_0.5 high-entropy alloys that the cooling rate affects the strength by inhibiting the precipitation of MoC, and the yield strength in the furnace-cooled state is 15% higher than that in the water-cooled state. Huang et al. [21] studied the influence of cooling rate on the phase stability of HfTaTiZr and HfNbTiZr high-entropy alloys and found that the former showed a strong dependence of microstructure and tensile properties on the cooling path, while the latter did not. The HfTaTiZr alloy retained the BCC lattice under water quenching but underwent complex phase decomposition during air cooling and furnace cooling. For composites, the effect of cooling methods is more complex: on the one hand, rapid cooling may inhibit the element diffusion between reinforcing phases (such as TiB₂) and the matrix, affecting the solid solution strengthening effect; on the other hand, slow cooling may promote uniform element distribution and enhance interface bonding. Zhang et al. [22] found that rapid cooling can greatly improve the strength of in situ TiC-reinforced CoCrFeNi high-entropy alloy matrix composites. However, in the current research on EHEAMCs, the correlation mechanism of cooling methods on element diffusion at the “reinforcing phase-matrix” interface, phase structure evolution, and mechanical properties is still unclear. Especially in the TiB₂@Ti-reinforced AlCoCrFeNi_2.1 system, there is no systematic understanding of how the diffusion behavior of Ti and B elements is regulated by the cooling rate and then affects the synergy between solid solution strengthening and interface strengthening.

Based on this, the present study takes 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs as the research object, focusing on the influence of cooling methods (FC, AC, and WC) on their microstructure and mechanical properties. The phase structure and element distribution are characterized by XRD, SEM, and EDS, and the mechanical properties are evaluated by microhardness testing and room-temperature compression experiments to reveal the mechanism of cooling rate on FCC/BCC phase separation, Ti/B element diffusion, and solid solution strengthening. The research aims to clarify the internal relationship of “cooling method-microstructure-mechanical properties”, provide a regulable process path for the performance optimization of EHEAMCs, and at the same time provide theoretical support for the design of high-performance structural materials in fields such as marine engineering.

2. Experimental Materials and Methods

This experiment took 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs prepared by powder metallurgy (sintering temperature is 1050 °C) as the research object. First, the composites were heat-treated in a box-type resistance furnace at 1000 °C for 12 h and then subjected to three cooling methods, that is, FC, AC, and WC, aiming to study the influence of different cooling methods on the microstructure and mechanical properties of the composites. After heat treatment, the three types of EHEAMCs were ground with SiC sandpaper and polished with diamond paste before microstructure analysis and microhardness testing. A Bruker D8 ADVANCE X-ray diffractometer (XRD, Billerica, MA, USA) and a Quanta 250 scanning electron microscope (SEM, FEI Company, Hillsboro, OR, USA) were used to analyze the microstructure and morphology of the EHEAMCs. The XRD test conditions were as follows: Cu-Kα was used for X-ray measurement with a wavelength of 1.54056 Å; the operating tube voltage and tube current were 40 kV and 40 mA, respectively; the scanning angle ranged from 20° to 90° (2θ) with a scanning speed of 5°/min and a scanning step of 0.02°. The micro-area composition of the EHEAMCs was analyzed using an energy dispersive X-ray spectrometer (EDS) attached to the SEM system. Microhardness was measured with an HV-1000SPTA microhardness tester (YIMA, Shenzhen, China) (F = 9.8 N, T = 30 s). Each sample was measured five times, and the average value was taken as the measured hardness. A universal testing machine (MTS810, MTS Systems Corporation, Eden Prairie, MN, USA) was used to test the compressive properties of the samples with uniaxial compression at a speed of 1 mm/min. Three samples were measured for each type of composite, and the average value was taken as the compressive property of the composite.

3. Results and Discussion

3.1. XRD Results

Figure 1 shows the XRD patterns of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs in the as-sintered state and under different cooling conditions (FC, AC, WC). The XRD patterns show that the composite materials in all states contain a face-centered cubic (FCC) solid solution phase and a body-centered cubic (BCC) solid solution phase. Based on the designed composition of the composite (containing 10 wt.% TiB₂@Ti reinforcing phase), it is inferred that TiB₂ phase and the Ti phase may exist in the system. However, their contents may be lower than the detection limit of XRD, so no identifiable characteristic peaks are formed in the patterns. The cooling method does not change the existing types of FCC and BCC phases but only affects the intensity and position of their diffraction peaks. From the overall diffraction characteristics (Figure 1a), the characteristic peaks of each phase in the sintered state are clear, indicating that a stable crystal structure has been formed after vacuum hot-pressing sintering at 1050 °C. Compared with the sintered state, there is no obvious shift in the positions of the diffraction peaks of each phase in the FC, AC, and WC states, which confirms that the cooling process does not cause phase transformation, but there are differences in peak intensity. There are certain differences in the diffraction peak intensities among the FC, AC, and WC states (FC is slightly higher than the sintered state, followed by AC, and WC changes relatively obviously). However, these differences are small and may be affected by factors such as surface preparation accuracy, sample thickness uniformity, or instrument noise. They cannot be directly attributed to the degree of atomic diffusion or differences in crystal orientation for the time being, and their internal correlation needs further verification.

A further observation of the 40–50° diffraction range (Figure 1b) reveals that the cooling rate has a significant impact on the separation degree of the diffraction peaks of the FCC and BCC phases. In the as-sintered state, the diffraction peaks of the two phases almost overlap, which is because atoms have sufficient time to diffuse after sintering, masking the lattice characteristics of the two phases. In the range of 40–50°, with the increase in cooling rate (from FC to AC to WC), the diffraction peaks of the FCC and BCC phases show a tendency of gradual separation. This phenomenon may be related to the fact that rapid cooling inhibits atomic migration, leading to a prominent difference in lattice parameters between the two phases, but this mechanism still needs further verification. During slow cooling, atoms can diffuse across phases and adjust their arrangement, leading to the weakening of the difference in lattice parameters between the two phases; however, during rapid cooling, atoms are difficult to diffuse, and the independence of the lattice characteristics of the FCC and BCC phases becomes prominent, eventually manifesting as the separation of diffraction peaks.

The diffraction peak of the FCC phase (111) crystal plane shifts toward smaller angles as the cooling rate increases. A similar phenomenon has also been reported in studies by Zhou [19], Yao [20], and others, yet no reasonable explanation has been provided. This paper hypothesizes that this phenomenon may be associated with the diffusion behavior of Ti elements: the atomic radius of Ti is larger than that of the principal element in the FCC phase, and its solid solution is likely to induce lattice expansion (an increase in lattice constant), which in turn causes the diffraction angle to shift toward smaller angles. The difference in the offset between slow cooling (FC) and fast cooling (WC) might be related to the retention state of Ti atoms within the FCC phase, though this hypothesis requires further verification through direct characterization of element distribution.

3.2. Microstructure

Figure 2 shows the microstructure morphology of the as-sintered TiB₂@Ti/AlCoCrFeNi_2.1 eutectic high-entropy alloy matrix composite. Low-magnification observation (Figure 2a) reveals that the composite consists of a light gray eutectic high-entropy alloy matrix (EHEA) and a dark black grain boundary zone (GB Zone). The eutectic high-entropy alloy matrix exhibits a typical lamellar structure, with numerous black particles distributed in the grain boundary zone. High-magnification observation (Figure 2b) shows that a dark gray solid solution zone (SS Zone) exists in the eutectic high-entropy alloy matrix adjacent to the grain boundary. Combined with the EDS analysis results (Figure 2e,f,

Table 1. EDS results of TiB₂@Ti/AlCoCrFeNi_2.1EHEAMCs under as-sintered and different cooling conditions (at.% cooling conditions (at.%)).

Sample	Point	Al	Co	Cr	Fe	Ni	Ti	B
As-Sintered	1	1.4	1.2	1.3	1.7	2.3	75.5	16.5
	2	16.5	16.5	3.4	9.1	38.0	16.3	0.2
	3	5.2	6.1	58.0	8.0	10.6	6.6	5.5
	4	5.5	19.1	14.3	20.9	33.6	6.6	0
FC	1	3.3	2.7	1.0	2.4	6.4	46.2	38.1
	2	13.7	15.9	11.2	16.7	30.8	11.3	0.3
	3	3.7	6.7	56.1	9.5	12.0	3.9	8.1
	4	6.9	19.7	12.8	20.3	32.9	7.4	0
AC	1	2.4	1.6	1.8	1.4	3.8	58.0	30.9
	2	15.0	14.3	4.8	10.3	35.3	17.5	2.8
	3	6.5	9.2	40.5	8.8	21.0	6.7	7.4
	4	5.8	20.1	13.2	21.4	33.5	6.0	0
WC	1	0.8	0.8	0.5	0.8	1.9	61.2	34.0
	2	11.6	11.9	5.6	10.2	25.9	26.2	8.6
	3	3.7	5.2	65.3	8.9	9.0	3.6	4.3
	4	4.7	20.1	14.4	22.0	34.2	4.6	0

) and existing research conclusions (Wani et al. [8] confirmed that the dark-colored region of the eutectic structure in AlCoCrFeNi_2.1 is an Al-Ni-rich BCC phase, and the light-colored region is a Co-Fe-Ni-rich FCC phase), it can be further clarified that the black particles (Point 1) in the grain boundary zone mainly contain Ti (75.5 at.%) and B (16.5 at.%), which should be TiB₂ phase and a small amount of Ti phase. The light white region (Point 4) in the matrix is mainly composed of Ni (33.6 at.%), Fe (20.9 at.%), and Co (19.1 at.%), corresponding to the FCC phase. The dark gray region (Point 3) has a Cr content of 58.0 at.% and contains B (5.5 at.%), corresponding to the BCC phase. The solid solution zone (Point 2) shows a mixed distribution of elements such as Ni (38.0 at.%), Al (16.5 at.%), and Co (16.5 at.%).

Figure 3, Figure 4 and Figure 5 show the SEM morphologies and EDS analysis spectra of the composite material under FC, AC, and WC states, respectively. Meanwhile, the EDS data of each region in the composite material are listed in Table 1. Overall, the three cooling methods did not significantly alter the basic composition of the microstructure (eutectic matrix, solid solution zone, and grain boundary zone) but had a significant impact on element distribution, a process driven by both atomic diffusion kinetics and thermodynamic affinity. Under slow cooling conditions (FC and AC), Ti and B atoms had sufficient time to diffuse from the grain boundary zone to the matrix. Due to the strong negative mixing enthalpy between Ti and Ni, Al (−35 to −28 kJ·mol⁻¹ [23]), Ti preferentially dissolved in the FCC phase (Point 4); while B tended to enrich in the BCC phase (Point 3) because of its negative mixing enthalpy with Cr (−31 kJ·mol⁻¹ [23]). This selective diffusion resulted in higher Ti content in the FCC phase (7.4 at.% in the FC state) and higher B content in the BCC phase (8.1 at.% in the FC state) compared to fast cooling conditions. In contrast, rapid cooling (WC) inhibited atomic migration by shortening the diffusion time window, confining Ti and B to the grain boundary zone (Ti: 61.2 at.%, B: 34.0 at.% in Point 1 of the WC state). This led to lower concentrations of Ti (4.6 at.% in the FCC phase) and B (4.3 at.% in the BCC phase) in the matrix, but a more uniform element distribution in the solid solution zone (Point 2 of the WC state) because limited diffusion prevented selective enrichment. Under all conditions, the diffusion pathway was dominated by grain boundary diffusion (with low activation energy); volume diffusion contributed to further homogenization only under slow cooling (FC), which also explained why the Ti/B concentrations in the matrix were higher at this time.

3.3. Mechanical Properties

3.3.1. Hardness

Figure 6 shows the microhardness distribution characteristics of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs in the as-sintered state and under different cooling methods (FC, AC, WC), with specific data listed in

Table 2. Microhardness of TiB₂@Ti/AlCoCrFeNi_2.1 eutectic high-entropy alloy matrix composite in as-sintered state and under different cooling methods (Matrix: Matrix; GB Zone: Grain Boundary Zone; Average: Average Value).

Status	Zone	Harness
As-Sintered		546.5 ± 22.3
FC	Matrix	601.0 ± 27.3
	GB Zone	653.1 ± 24.8
	Average	627.0 ± 26.1
AC	Matrix	551.3 ± 11.8
	GB Zone	591.2 ± 34.3
	Average	571.3 ± 23.1
WC	Matrix	511.6 ± 18.9
	GB Zone	574.7 ± 51.9
	Average	543.2 ± 35.4

. The results indicate that the cooling method significantly affects the hardness performance of the material by regulating element diffusion behavior, showing a regular change that “the lower the cooling rate, the higher the hardness”. In terms of overall hardness level, after heat treatment at 1000 °C for 12 h, the average hardness of the three cooling states shows differentiated characteristics. The FC state has the highest average hardness, reaching 627.0 ± 26.1 HV; the AC state comes next at 571.3 ± 23.1 HV; and the WC state is the lowest at 543.2 ± 35.4 HV. It is worth noting that although the WC state has the lowest hardness, it is still slightly higher than the as-sintered state (546.5 ± 22.3 HV), indicating that rapid cooling after heat treatment can still retain part of the strengthening effect. Analysis from the perspective of micro-area hardness distribution (matrix and grain boundary zone) shows that the influence of cooling rate is more significant. In the FC state, the hardness of the grain boundary zone (653.1 ± 24.8 HV) is higher than that of the matrix (601.0 ± 27.3 HV), which is related to the sufficient diffusion of Ti and B elements from the grain boundary to the matrix under slow cooling conditions. The concentration of residual TiB₂ particles (a high-hardness phase) in the grain boundary zone is relatively high (the atomic fractions of Ti and B are 46.2 at.% and 38.1 at.%, respectively, in Table 1), and at the same time, the matrix has enhanced lattice distortion caused by the solid solution of Ti elements, forming a synergistic effect of “grain boundary particle strengthening + matrix solid solution strengthening”. With the increase in cooling rate, the hardness difference between the grain boundary and the matrix gradually narrows. In the AC state, the hardness of the grain boundary zone (591.2 ± 34.3 HV) decreases by about 9.5% compared with the FC state, and the matrix hardness (551.3 ± 11.8 HV) decreases by about 8.3%. In the WC state, the hardness of the grain boundary zone (574.7 ± 51.9 HV) decreases by about 12.0% compared with the FC state, and the matrix hardness (511.6 ± 18.9 HV) decreases by about 14.9%. The core reason for this phenomenon is that rapid cooling inhibits the diffusion of Ti and B elements (the Ti content in the FCC phase of the WC state matrix is only 4.6 at.% in Table 1, which is much lower than 7.4 at.% in the FC state), resulting in the weakening of the solid solution strengthening effect. At the same time, the distribution uniformity of TiB₂ particles in the grain boundary zone decreases (the degree of element segregation is reduced), and the particle strengthening effect is weakened. In summary, the cooling rate affects the synergistic effect of solid solution strengthening and particle strengthening by regulating the diffusion behavior of Ti and B elements. Slow cooling (FC) promotes sufficient diffusion of elements, making the strengthening mechanism more significant; rapid cooling (WC) inhibits diffusion, weakening the strengthening effect, but it is still better than the as-sintered state without heat treatment.

3.3.2. Room-Temperature Compression Property

Figure 7 shows the room-temperature compressive stress–strain curves of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs in the as-sintered state and under different cooling methods (FC, AC, WC), and

Table 3. Room-temperature compressive properties of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs in as-sintered state and under different cooling methods.

Status	Zone	Harness
As-Sintered		546.5 ± 22.3
FC	Matrix	601.0 ± 27.3
	GB Zone	653.1 ± 24.8
	Average	627.0 ± 26.1
AC	Matrix	551.3 ± 11.8
	GB Zone	591.2 ± 34.3
	Average	571.3 ± 23.1
WC	Matrix	511.6 ± 18.9
	GB Zone	574.7 ± 51.9
	Average	543.2 ± 35.4

lists the corresponding compressive mechanical property parameters (yield strength σ₀.₂, fracture strength σmax, fracture strain εmax). The results show that the cooling rate has a significant impact on the compressive properties of the composites, generally presenting a rule that “mechanical properties improve as the cooling rate decreases”, and the properties after heat treatment are all better than those in the as-sintered state. In terms of specific data, the FC state (with the slowest cooling rate) has the optimal comprehensive mechanical properties. Its yield strength reaches 1574 MPa, which is approximately 14.4% higher than that of the as-sintered state (1376 MPa); the fracture strength reaches 2824 MPa, an increase of about 18.2% compared with the as-sintered state (2388 MPa); the fracture strain reaches 24.2%, which is about 18.0% higher than that of the as-sintered state (20.5%). The AC state has the second-best properties, with a yield strength of 1481 MPa, a fracture strength of 2758 MPa, and a fracture strain of 23.8%, which are approximately 5.9%, 2.3%, and 1.7% lower than those of the FC state, respectively. The WC state (with the fastest cooling rate) has the lowest properties but is still better than the as-sintered state. Its yield strength is 1401 MPa (about 1.8% higher than that of the as-sintered state), the fracture strength is 2636 MPa (approximately 10.4% higher than that of the as-sintered state), and the fracture strain is 23.0% (about 12.2% higher than that of the as-sintered state).

The compressive properties of the composites show a clear dependence on cooling rate, with mechanical performance improving as the cooling rate decreases—consistent with the microhardness trend but with distinct characteristics in strength-plasticity synergy. The FC state exhibits the optimal comprehensive compressive properties, with yield strength (1574 MPa), fracture strength (2824 MPa), and fracture strain (24.2%) all significantly higher than those of the as-sintered state. This is attributed to the sufficient diffusion of Ti and B under slow cooling, which enhances solid solution strengthening (as discussed in Section 3.3.1) and optimizes element distribution between the FCC and BCC phases. This synergy avoids local stress concentration during compression, enabling simultaneous improvement in strength and plasticity. As the cooling rate increases (AC→WC), the compressive strength decreases gradually: the AC state shows a 5.9% reduction in yield strength and a 2.3% reduction in fracture strength compared to FC, while the WC state exhibits further reductions (11.0% in yield strength and 6.6% in fracture strength relative to FC). However, the WC state retains better plasticity (fracture strain: 23.0%) than the as-sintered state, which is related to the more uniform distribution of Ti and B in the matrix under rapid cooling. This uniformity reduces intra-phase composition fluctuations and local stress concentration points, inhibiting crack initiation and propagation during compression. These results confirm that cooling rate regulates compressive behavior primarily through controlling the extent of Ti/B diffusion and the resulting solid solution strengthening effect. Slow cooling (FC) achieves the best strength-plasticity balance, while rapid cooling (WC) maintains moderate performance due to improved element uniformity.

4. Conclusions

The present study focused on 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs prepared by powder metallurgy (sintering temperature: 1050 °C). The composites were first subjected to heat treatment at 1000 °C for 12 h and then cooled by three methods: furnace cooling (FC), air cooling (AC), and water cooling (WC). The microstructure and phase structure were analyzed by XRD, SEM, and EDS, and the mechanical properties were evaluated by microhardness testing and room-temperature compression experiments. The main conclusions are as follows:

(1) In the as-sintered state and under different cooling methods, the composites are all composed of FCC phase, BCC phase, TiB₂, and Ti phase. The cooling methods do not change the phase types but affect the intensity and position of diffraction peaks. With the increase in cooling rate, the diffraction peaks of the FCC and BCC phases gradually separate from the overlapping state in the as-sintered state, and the diffraction peak of FCC (111) crystal plane shifts to a lower angle (due to the increase in lattice constant caused by Ti element diffusion). The diffraction peak intensity decreases with the increase in cooling rate. The peak intensity of the FC state is slightly higher than that of the as-sintered state because of sufficient atomic diffusion, and the peak intensity of the WC state changes most significantly due to the inhibition of atomic diffusion.

(2) In the as-sintered state and under the three cooling conditions, the composites are all composed of a eutectic high-entropy alloy matrix, a solid solution zone, and a grain boundary zone. The cooling rate has little effect on the morphology but has an obvious effect on the element distribution. At low cooling rates (FC, AC), Ti and B elements diffuse fully from the grain boundary to the matrix, resulting in higher concentrations of both in the matrix. Moreover, Ti preferentially dissolves in the FCC phase, and B preferentially dissolves in the BCC phase. During rapid cooling (WC), element diffusion is inhibited, so the concentrations of Ti and B in the matrix are low (Ti in FCC phase: 4.6 at.%, B in BCC phase: 4.3 at.%), but the element distribution is more uniform.

(3) The mechanical properties of the composites decrease with the increase in cooling rate. The FC state has the best performance, and the WC state has the worst performance but is still better than the as-sintered state. The average hardness of the FC state is 627.0 ± 26.1 HV, with a yield strength of 1574 MPa, a fracture strength of 2824 MPa, and a fracture strain of 24.2%. The average hardness of the WC state is 543.2 ± 35.4 HV, with a yield strength of 1401 MPa, a fracture strength of 2636 MPa, and a fracture strain of 23.0%. The main strengthening mechanism is solid solution strengthening. Slow cooling promotes the sufficient diffusion of Ti and B elements into the matrix, enhances lattice distortion, and improves the strength and plasticity. In the future, we will further investigate the service performance of this composite material in extreme environments (such as high temperatures and corrosive media) and expand research on its application in the field of aerospace precision components.