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

Abstract

This study systematically investigates the effects of heat treatment at 800–1000 °C on the microstructure and mechanical properties of 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 eutectic high-entropy alloy matrix composites (EHEAMCs) prepared by vacuum hot-pressing sintering. The results show that the materials consist of FCC, BCC, TiB₂, and Ti phases, with a preferred orientation of the (111) crystal plane of the FCC phase. As the temperature increases, the diffraction peak of the BCC phase separates from the main FCC peak and its intensity increases, while the diffraction peak positions of the FCC and BCC phases shift at small angles. This is attributed to the diffusion of TiB₂@Ti from the grain boundaries into the matrix, where the Ti solid solution increases the lattice constant of the FCC phase. Microstructural observations reveal that the eutectic region transforms from lamellar to island-like structures, and the solid solution zone narrows. With increasing temperature, the Ti concentration in the solid solution zone increases, while the contents of elements such as Ni decrease. Element diffusion is influenced by binary mixing enthalpy, with Ti and B tending to solidify in the FCC and BCC phase regions, respectively. The mechanical properties improve with increasing temperature. At 1000 °C, the average hardness is 579.2 HV, the yield strength is 1294 MPa, the fracture strength is 2385 MPa, and the fracture strain is 19.4%, representing improvements of 35.5% and 24.9% compared to the as-sintered state, respectively, without loss of plasticity. The strengthening mechanisms include enhanced solid solution strengthening due to the diffusion of Ti and TiB₂, improved grain boundary strength due to the diffusion of alloy elements to the grain boundaries, and synergistic optimization of strength and plasticity.

1. Introduction

As a new metal material system breaking through traditional alloy design concepts, high-entropy alloys (HEAs) have become a research hotspot in the field of metal materials in recent years due to their unique multi-principal-element synergistic effects and excellent mechanical properties and functional characteristics [1,2,3]. In particular, eutectic high-entropy alloys (EHEAs), which achieve a good synergistic effect of strength and plasticity through the coupling of multiphase structures and lamellar microstructures, provide an important direction for the development of new structural materials [4]. The AlCoCrFeNi_2.1 system, as a typical representative of EHEAs, exhibits a lamellar eutectic structure of FCC and B2/BCC phases in its as-cast state, with a yield strength as high as 1.2 GPa and a fracture elongation of over 15% [5], showing broad application prospects in aerospace high-temperature components, ship corrosion-resistant structural components, and other fields.

To further improve the comprehensive properties of EHEAs, researchers have mainly explored two directions: alloying modification and composite strengthening. In terms of alloying, introducing large atomic radius elements (such as Ti, Nb, Mo, etc.) to produce lattice distortion strengthening effects has become an effective strategy. Kratochvíl et al. [6] found that adding Ti elements can increase the hardness of CoCrFeNi alloys to 968 HV and the compressive yield strength to 2157 MPa, which is attributed to the dual strengthening mechanisms of lattice distortion caused by Ti atoms and the precipitation of Laves phases. He et al. [7] successfully constructed a nanolamellar eutectic structure in CoCrFeNiNb_0.5 alloys, achieving a compressive strength of 2.3 GPa while maintaining a compressive strain of 23.6%, confirming the optimization effect of multi-scale structural regulation on mechanical properties. However, the single alloying strategy often leads to plastic loss while improving strength. For example, Shun et al. [8] reported that the fracture strain of CoCrFeNiMo_x alloys dropped sharply to 21% when the Mo content reached 0.5.

Composite strengthening provides a new idea to break through the strength–plasticity inversion relationship. Constructing high-entropy alloy matrix composites (HEAMCs) by introducing ceramic reinforcing phases can combine the ductility of the matrix with the high modulus of the reinforcing phases. For example, Yim et al. [9] added 5 vol.% TiC particles to a CoCrFeMnNi matrix, increasing the yield strength by 37.5% while maintaining the original plastic level. Rogal et al. [10] used SiC nanoparticles to reinforce HEAMCs, achieving a room-temperature compressive strength of 1480 MPa, which is 25.4% higher than that of the matrix. Typical microstructural features of current HEAMCs include uniform distribution of ceramic phases in the high-entropy alloy matrix, and the formation of interfacial reaction layers to improve compatibility. For example, Zhou et al. [11] formed M₇C₃-type carbide layers in the WC/FeCoCrNi system, reducing the wear rate by 40%. Wang et al. [12] prepared xZrB₂/AlCoCrFeNi (x = 0, 5, 10, 15 wt%)-based HEAMCs by spark plasma sintering (SPS), whose phase structure experienced a transition from (FCC + BCC + B2) to (FCC + BSC + B2 + Laves), with a maximum compressive strength of 2071 MPa and a maximum hardness of 1222 HV. In research on AlCoCrFeNi_2.1 EHEAMCs, the TiB₂/AlCoCrFeNi_2.1 EHEAMCs prepared by Han et al. [13] had a microhardness of 780 HV and a compressive strength of 2.5 GPa, and the strengthening mechanism was mainly attributed to the pinning effect of TiB₂ particles and fine-grain strengthening. Jiang et al. [14] found that with the increase in NbC content, the hardness of NbC/AlCoCrFeNi_2.1 EHEAMCs ranged from 660 to 700 HV, and the average friction coefficient decreased. Liu et al. [15] found that the interface of WC/AlCoCrFeNi_2.1 EHEAMCs reacted to form M₇C₃ and M₂₃C (M = Cr, W) phases, forming a reaction layer with a thickness of 1.4–1.9 μm, which significantly improved the hardness of the composite. The performance advantages of these HEAMCs include the synergy of high strength and good wear resistance, but in traditional designs, ceramic phases easily lead to plastic loss, the interfacial structure stability is poor, and the interfacial element diffusion mechanism is not clear.

Heat treatment, as a key post-processing technology for regulating material microstructure, has unique advantages in optimizing the properties of HEAs. Sun et al. [16] found that aging treatment at 500 °C induced the precipitation of the σ phase in Cr_0.8FeMn_1.3Ni_1.3 alloys, increasing the hardness to 450 HV. Yin et al. [17] studied Fe₄₅Ni₂₅Cr₂₅Mo₅ alloys and showed that the σ phase precipitation after aging at 900 °C for 48 h increased the compressive strength by 85%. Shafiei et al. [18] studied Al₁₄Co₄₁Cr₁₅Fe₁₀Ni₂₀ HEAs and found that both as-cast and homogenized (1200 °C/20 h) alloys had a dual-phase structure of dendritic FCC solid solution and interdendritic B2 (BCC) phase. After aging treatment at 750–1000 °C, precipitates were uniformly distributed, and with increasing temperature, the precipitates coarsened, and the hardness decreased due to over-aging. For AlCoCrFeNi_2.1 EHEAs, Wani et al. [19] found that after 90% cold rolling and heat treatment at 800–1200 °C, the cold rolling disordered the L1₂ phase while the B2 phase remained ordered, and the structure transformed from lamellar to equiaxed grains. With increasing temperature, the growth of the two phases hinders each other, leading to grain refinement. Bhattacharjee et al. [20] improved the comprehensive mechanical properties of AlCoCrFeNi_2.1 EHEAs after low-temperature rolling and heat treatment at 800 °C. Reddy et al. [21] studied AlCoCrFeNi_2.1 EHEAs after warm rolling (90%) at 400 °C and 750 °C and heat treatment, finding that the L1₂ phase was disordered, part of the FCC phase remained ordered after warm rolling at 600 °C, and with increasing rolling and heat treatment temperatures, the hardness and yield strength of the alloys decreased, while the elongation increased. The comprehensive properties were better after warm rolling at 750 °C and heat treatment at 1000 °C for 1 h. Shi et al. [22] performed cold rolling (84–86%) and heat treatment with water quenching at different temperatures on AlCoCrFeNi_2.1 EHEAs, obtaining a dual-phase heterogeneous eutectic structure. With increasing temperature, the yield strength decreased, and the elongation increased. These studies show that reasonable rolling and heat treatment can effectively regulate the mechanical properties of AlCoCrFeNi_2.1 EHEAs. Wang et al. [23] found that aging AlCoCrFeNi_2.1 EHEAs prepared by laser melting at 500–700 °C broke the B₂ network and induced γ′ precipitation strengthening; aging at 800 °C dissolved the B₂ network and γ′, reducing the strength while increasing the plasticity. For HEAMCs, Yin et al. [24] in situ synthesized FeNiCrMoTiC HEAMCs with an FCC matrix, carbides, and eutectic structure. After aging at 800 °C for 96 h, the precipitation strengthening of intermetallic compounds improved hardness and strength, and the peak-aged state exhibited wear resistance, exceeding high-performance high-chromium cast iron. The unique structure inhibited severe delamination and crack propagation, and the welding of spalled phases strengthened the matrix, giving it both oxidation and corrosion resistance. However, existing research has mostly focused on the heat treatment effects of single alloy systems, and there is a significant lack of research on composite systems. In particular, under the condition of ceramic reinforcing phases, the element diffusion behavior and interfacial evolution law, and their influence mechanisms on mechanical properties induced by heat treatment, remain unclear [25].

It is worth noting that existing studies have paid insufficient attention to the heat treatment effects in the presence of ceramic reinforcing phases. Most of the work has focused on the heat treatment regulation of single-phase high-entropy alloys [16,17,18,19,20,21,22,23,24], while there is a lack of systematic understanding of the element diffusion behavior and interfacial reaction mechanisms, and their effects on the strength–ductility synergy induced by heat treatment in EHEAMCs. For example, Wang et al. [23] studied the aging behavior of AlCoCrFeNi_2.1 eutectic alloys, but did not consider the role of reinforcing phases, and although Han et al. [13] prepared TiB₂-reinforced composites to achieve high strength, this was accompanied by significant plastic loss. In addition, traditional composite material design mostly relies on the trial-and-error method to add reinforcing phases [15], lacking a microstructural regulation strategy based on thermodynamic principles. This study takes 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC as the research object, systematically studies the effects of heat treatment at 800–1000 °C on their microstructure and mechanical properties, and provides a new path for breaking through the strength–ductility synergy bottleneck of high-entropy alloy matrix composites by revealing the internal correlation among element diffusion, interfacial evolution, and performance optimization. The composite has the potential to replace traditional nickel-based alloys or metal matrix composites in aerospace high-temperature components due to its excellent high-temperature mechanical stability.

2. Experimental Materials and Methods

The raw materials used in the experiment were AlCoCrFeNi_2.1, TiB₂, and Ti powders, with average particle sizes of 85 μm, 3 μm, and 3 μm, respectively, and a purity of over 99%. The chemical composition was proportioned according to the stoichiometric ratio of 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC. The preparation of the composite is divided into two steps: first, the low-energy ball milling preparation of TiB₂@Ti powder and TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC powder. The TiB₂@Ti powder and TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC powder are prepared by low-energy ball milling using a horizontal planetary ball mill [26]. The ball milling tank and grinding balls are made of 304 stainless steel, with a ball-to-powder ratio of 5:1, a rotation speed of 150 rpm, Ar protection, and anhydrous ethanol as the process control agent. When preparing TiB₂@Ti powder, the molar ratio of Ti powder to TiB₂ powder is 1:1, and the ball milling time is 3 h. The ball-milled TiB₂@Ti is then mixed with AlCoCrFeNi_2.1 EHEA powder and ball-milled for 5 h. Second, the bulk TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC is prepared at 1000 °C using a QSH-ZR-55 type vacuum hot-pressing sintering furnace (Shanghai Quanshuo Electric Furnace Co., Ltd., Shanghai, China). During the hot-pressing sintering process, the background vacuum is 1 × 10⁻³ Pa, the heating rate is 15 °C/min, the sintering pressure is 30 MPa, the heat preservation time is 2 h, and after cooling with the furnace, cylindrical samples with the size of ϕ30 mm × 6 mm are obtained. The TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMC was then heat-treated in a box-type resistance furnace at 800 °C, 900 °C, and 1000 °C for 12 h of holding time, followed by furnace cooling. After heat treatment, the composites were ground with SiC sandpaper and polished with diamond polishing paste for microstructural analysis and microhardness testing. Compression specimens with dimensions of ϕ3 mm × 6 mm were also heat-treated simultaneously.

The microstructure and morphology of the EHEAMCs were analyzed using a Bruker D8 ADVANCE X-ray diffractometer (XRD) (Billerica, MA, USA) and a Quanta 250 scanning electron microscope (SEM) (Thermo Fisher Scientific, Waltham, MA, USA). The XRD test conditions were as follows: Cu-Kα X-ray measurement with a wavelength of 1.54056 Å, operating tube voltage and current of 40 kV and 40 mA, respectively, scanning angle of 20–90° (2θ), scanning speed of 5°/min, and scanning step size of 0.02°. The micro-area composition of the EHEAMCs was analyzed using an energy-dispersive X-ray spectrometer (EDS) integrated with the SEM system. The microhardness of the EHEAMCs was tested using an HV-1000SPTA microhardness tester (YIMA, Shenzhen, China) under the following experimental conditions: a load of 98 N and a holding time of 30 s. The microhardness of the matrix and grain boundaries of each specimen was measured at least five times, and the average value was taken as the evaluated hardness of the specimen. The compression properties of the composites were tested using an MTS810 universal testing machine (MTS Systems Corporation, Eden Prairie, MN, USA) in accordance with the ASTM E9-19 standard [27] test method for compression testing of metallic materials at room temperature. During the compression test, the loading rate was 2 × 10⁻⁴/s, and three specimens of each composite were tested. The average value of their performance indicators was taken as the compression performance of the composite.

3. Results and Discussion

3.1. Phase Structure

Figure 1 shows the XRD results of the TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs under different heat treatment temperatures. It can be seen that the phase structure of the composites in the as-sintered state and after heat treatment at different temperatures consists of FCC, BCC, as well as TiB₂ and Ti phases. Among them, the (111) crystal plane of the FCC phase exhibits preferred orientation (2θ ≈ 44.0°), indicating that the composite has a relatively stable phase structure under different heat treatment conditions. With the increase in heat treatment temperature, the diffraction peak of the BCC phase (2θ ≈ 44.8°) gradually separates from the main FCC peak, and its intensity increases (Figure 1b). This is mainly because the increase in heat treatment temperature provides sufficient energy for atomic diffusion, promoting the redistribution of elements and phase transformation in the alloy.

In addition, the diffraction peak positions of the FCC and BCC phases shift toward smaller angles with the increase in heat treatment temperature. As shown in

Table 1. Lattice constants and crystallite sizes of the FCC and BCC phases in the TiB₂@Ti/AlCoCrFeNi_2.1 composites at different heat treatment temperatures.

State	Lattice Constant (nm)		Crystallite Size (nm)
	FCC	BCC	FCC	BCC
As-Sintered	0.3562	0.2862	32.2	27.6
800 °C	0.3569	0.2869	60.9	50.2
900 °C	0.3572	0.2870	143.4	53.9
1000 °C	0.3580	0.2885	107.6	123.1

, heat treatment temperature significantly regulates the lattice constants and crystallite sizes of the FCC and BCC phases in the TiB₂@Ti/AlCoCrFeNi_2.1 composites. The lattice constant of the FCC phase increases gradually from 0.3562 nm in the as-sintered state to 0.3580 nm at 1000 °C, with an increase of approximately 0.5%. This is attributed to the diffusion and solid solution of the Ti elements, with a larger atomic radius (147 pm) from TiB₂@Ti at the grain boundaries into the FCC phase, causing strong lattice distortion. The diffraction peak shift of the FCC phase to a smaller angle further confirms lattice expansion. This is consistent with the phenomenon observed by Kratochvíl et al. [6] in the CoCrFeNiTi_x alloy system, where lattice parameter increase and diffraction peak shift were caused by Ti element diffusion.

The lattice constant of the BCC phase increases from 0.2862 nm to 0.2885 nm, with an increase of approximately 0.8%, which is due to the easy solid solution of B elements into the BCC phase because of the binary mixing enthalpy of −31 kJ·mol⁻¹ between B and Cr. The cumulative effect of the interstitial solid solution leads to lattice expansion. In terms of crystallite size, the FCC phase first increases from 32.2 nm to 143.4 nm at 900 °C and then decreases to 107.6 nm at 1000 °C. This is because medium and low temperatures promote normal grain growth, while high temperatures cause an increase in lattice distortion energy due to the Ti solid solution and the pinning effect of the TiB₂ particles at the grain boundaries, which inhibits excessive grain growth. The BCC phase continuously increases from 27.6 nm to 123.1 nm, with an increase of 345%. This is because the BCC phase (rich in Al-Ni) has high thermodynamic stability at high temperatures, and the solid solution of B elements strengthens the lattice. In addition, the diffusion of elements from the FCC phase to the grain boundaries makes the composition of the BCC phase more uniform, promoting abnormal grain growth. The changes in lattice constants and crystallite sizes are essentially the result of the coupling between element diffusion (driven by binary mixing enthalpy) and phase structure evolution. Medium and low temperatures promote atomic diffusion and grain growth, while high temperatures coexist with recrystallization of the FCC phase and abnormal growth of the BCC phase, providing a theoretical basis for optimizing the microstructure and improving the mechanical properties through heat treatment.

3.2. Microstructure

Figure 2 shows the SEM images of the TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs after heat treatment at different temperatures. The microstructure of the composites exhibits typical multiphase characteristics, with the matrix phase distributed continuously, while the TiB₂ and Ti particles serve as reinforcing phases dispersed in the matrix. In the as-sintered state, the TiB₂ particles vary in size, and partial particle agglomeration occurs, which may be due to interactions between the particles during preparation and limitations of the process conditions. With the increase in heat treatment temperature, the agglomerated TiB₂ particles gradually disperse, and the particle size is refined to some extent. This phenomenon may be attributed to intensified atomic diffusion at high temperatures, which allows originally agglomerated particles to redistribute. Meanwhile, thermodynamic processes during heat treatment, such as the reverse effect of the Ostwald ripening mechanism, inhibit the growth trend of large particles by promoting the dissolution of small particles, thereby leading to particle size refinement. As shown in Figure 2a, the matrix of the as-sintered composite has a lamellar structure, with numerous black particles at the grain boundaries. The gray area between the grain boundaries and the matrix is the solid solution zone (SS Zone) of alloying elements. In the matrix, the dark structure is the Al-Ni-rich BCC phase, and the light structure is the Co-Fe-Ni-rich FCC phase [28]. After heat treatment at 800 °C, the morphology of part of the eutectic region in the composite transforms from lamellar to island-like, and the solid solution zone narrows (Figure 2b). When the temperature rises to 1000 °C, most of the lamellar eutectic regions transform into island-like regions, the solid solution zone becomes smaller, and several micron TiB₂ particles are distributed at the grain boundaries (Figure 2d).

According to the EDS results listed in

Table 2. EDS analysis results (at.%) of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs under different heat treatment temperatures.

Status	Point	Al	Co	Cr	Fe	Ni	Ti	B
As-Sintered	1	0.4	0.4	0.3	0.3	0.7	65.8	32.1
	2	21.4	10.6	9.3	8.9	31.2	18.2	0.4
	3	27.5	15.3	10.2	12.3	25.1	8.8	0.8
	4	13.8	18.9	19.1	15.6	30.2	2.1	0.3
800 °C	5	0.3	0.2	0.4	0.2	0.6	68.0	30.3
	6	23.2	11.3	1.9	5.0	38.1	20.1	0.4
	7	7.9	14.1	25.2	10.5	30.2	11.5	0.6
	8	3.8	21.5	18.0	24.8	29.1	2.5	0.3
900 °C	9	0.7	0.4	0.3	0.4	0.7	67.9	29.6
	10	18.0	6.9	13.1	10.6	20.7	23.9	6.8
	11	10.4	15.8	25.3	19.1	15.4	13.6	0.4
	12	11.2	16.0	15.4	18.0	35.5	3.7	0.2
1000 °C	13	0.3	0.3	0.4	0.8	0.6	64.1	33.5
	14	13.5	12.1	12.8	8.4	27.8	25.1	0.3
	15	6.4	8.9	42.4	9.8	19.6	5.2	7.7
	16	11.2	10.4	26.5	8.6	25.3	13.2	4.8

, the main component elements of the black blocky substances in the GB zone (points 1, 5, 9, and 13 in Figure 2) are Ti and B, with the Ti concentration as high as over 64 at.%. In the solid solution zone (points 2, 6, 8, and 12 in Figure 2), the main component elements are Ni, Ti, Al, etc. With the increase in heat treatment temperature, the Ti concentration in the solid solution zone increases, while the contents of Ni and other alloying elements decrease. At higher heat treatment temperatures, the main component elements of the dark gray island-like structure in the eutectic high-entropy alloy matrix are Cr, Ni, B, etc., and the main component elements of the light white area are Ni, Fe, Co, etc. Combining with the XRD analysis results, it can be inferred that the black blocky substances in the GB zone are likely aggregates of TiB₂ and Ti, the tiny particles may be TiB₂, the light white area in the EHEA matrix is the Ni-Fe-Co-rich FCC phase, and the dark gray island-like area is the Cr-Al-Ni-B-rich BCC phase. Compared with the as-sintered state, after high-temperature heat treatment at 1000 °C for 12 h, the mutual diffusion of component elements causes elements such as Ni, Fe, Al, and Co in the matrix to diffuse toward the grain boundaries, while the Ti and B elements at the grain boundaries diffuse into the matrix, resulting in an insignificant solid solution zone. In addition, the Ti elements tend to solidify in the FCC phase zone, while the B elements tend to solidify in the BCC phase zone, which is related to the binary mixing enthalpy between alloying elements. The binary mixing enthalpies of Ti-Ni, Ti-Al, and Ti-Co are −35, −30, and −28 kJ·mol⁻¹, respectively. Compared with other elements, they have larger negative enthalpies, so Ni, Al, and Co are more likely to diffuse toward the grain boundaries, resulting in higher Ni content at the grain boundaries. The binary mixing enthalpies of B-Cr and B-Ni are −31 and −24 kJ·mol⁻¹, respectively, so they are prone to solid solution in the Cr-rich BCC phase [29]. Notably, the microstructural evolution of the grain boundary (GB) region after heat treatment at 800 °C is particularly significant: the reduction in black lumps and the shrinkage of the solid solution zone are attributed to the interfacial dissociation of the TiB₂@Ti core-shell phase and directional element diffusion. The Ti atoms solidify into the matrix due to the low mixing enthalpy with the FCC phase, while the B atoms migrate to the BCC phase. Meanwhile, the reverse diffusion of matrix elements to the grain boundaries reconstructs the interfacial composition gradient, ultimately leading to the “suppression” phenomenon of the GB region observed by SEM (Figure 2b). This behavior is mutually verified by the collaborative changes in the lattice constants of the FCC/BCC phases (Table 1) and the element concentration distribution (Table 2).

3.3. Mechanical Properties

3.3.1. Hardness

The microhardness of the TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs at different heat treatment temperatures is shown in Figure 3 and

Table 3. Microhardness of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs at different heat treatment temperatures.

Status	Zone	Harness
As-Sintered		467.3 ± 21.2
800 °C	Matrix	479.4 ± 20.0
	GB Zone	567.5 ± 19.4
	Average	523.4 ± 19.7
900 °C	Matrix	520.0 ± 36.3
	GB Zone	559.9 ± 49.5
	Average	540.0 ± 42.9
1000 °C	Matrix	539.9 ± 29.6
	GB Zone	618.4 ± 59.8
	Average	579.2 ± 44.7

. The microhardness of the composites exhibits a significant increasing trend with the rise in heat treatment temperature. In the as-sintered state, the average hardness of the composite is 467.3 ± 21.2 HV. At this stage, the agglomeration of TiB₂ particles and uneven element distribution limit the strengthening effect. When the heat treatment temperature rises to 800 °C, the average hardness increases to 523.4 ± 19.7 HV, representing a 12% increase compared to the as-sintered state. The hardness of the matrix and grain boundaries increases to 479.4 HV and 567.5 HV, respectively, indicating the initial manifestation of local strengthening effects. Further increasing the temperature to 1000 °C, the average hardness reaches 579.2 ± 44.7 HV, a 23.9% increase compared to the as-sintered state. Notably, the grain boundary hardness exceeds 600 HV (618.4 ± 59.8 HV), and the matrix hardness increases to 539.9 HV, fully demonstrating the comprehensive strengthening effect. The hardness improvement exhibits a “gradient enhancement driven by temperature” characteristic.

Compared with the TiB₂/AlCoCrFeNi_2.1 composite prepared by Han et al. [13] (relying solely on particle pinning with a hardness of 780 HV), this study introduces “solid solution + grain boundary strengthening” through heat treatment. At the same TiB₂ content (10 wt.%), the hardness approaches 600 HV, with a strengthening efficiency improvement of over 30%. Traditional particle-reinforced composites rely on a single strengthening pathway, whereas this study activates multiple mechanisms through temperature regulation, avoiding the interfacial defects caused by reinforcing phase agglomeration and demonstrating a more balanced strengthening effect.

3.3.2. Room-Temperature Compression Property

Figure 4 shows the compressive stress–strain curves of the TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs at different heat treatment temperatures, and

Table 4. Room-temperature compressive properties of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs at different heat treatment temperatures.

Status	σ0.2/MPa	σmax/MPa	εmax/%
As-Sintered	955 ± 47	1910 ± 91	18.6 ± 0.9
800 °C	1023 ± 52	2206 ± 102	21.2 ± 0.7
900 °C	1137 ± 58	2219 ± 93	19.3 ± 0.7
1000 °C	1294 ± 61	2385 ± 111	19.4 ± 0.6

lists the relevant compressive mechanical property parameters. The compressive mechanical properties of the TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs show significant optimization with the increase in heat treatment temperature. In the as-sintered state, the yield strength, ultimate compressive strength, and fracture strain of the composite are 955 MPa, 1910 MPa, and 18.6%, respectively. At this stage, the agglomeration of TiB₂ particles and uneven element distribution make the grain boundaries weak regions for crack initiation. When the heat treatment temperature rises to 800 °C, the yield strength increases to 1023 MPa (an increase of 7.1%), the ultimate compressive strength reaches 2206 MPa (an increase of 15.5%), and the fracture strain increases to 21.2% (an increase of 14%), indicating that the initial stage of diffusion promotes both strength and plasticity through microstructural optimization. Further increasing the temperature to 1000 °C, the compressive properties of the composite reach their peak values: the yield strength, ultimate compressive strength, and fracture strain reach 1294 MPa (a 35.5% increase compared to the as-sintered state), 2385 MPa (a 24.9% increase), and 19.4%, respectively, achieving a synergistic improvement in strength and plasticity and breaking through the traditional bottleneck of the “strength–plasticity trade-off” in metallic materials.

Compared with similar HEAMCs, this system exhibits unique performance advantages. For example, although the TiB₂/AlCoCrFeNi_2.1 composite prepared by Han et al. [13] has a microhardness of 780 HV and a compressive strength of 2.5 GPa, its strengthening mechanism mainly relies on the pinning effect of the TiB₂ particles, with a fracture strain of only 8–10%, which is significantly lower in plasticity than the system in this study. Yim et al. [9] prepared a composite by adding 5 vol.% TiC particles to a CoCrFeMnNi matrix, which increased the yield strength by 37.5%; however, the fracture strain data were not specified, and the single-particle strengthening pathway is easily limited by the quality of interfacial bonding. The SiC nanoparticle-reinforced HEAMCs prepared by Rogal et al. [10] have a room-temperature compressive strength of 1480 MPa, a 25.4% increase over the matrix, but plasticity was not mentioned. In this current study, multiple mechanisms activated by heat treatment—“solid solution strengthening + grain boundary strengthening + particle dispersion strengthening + eutectic structure optimization”—not only significantly improve yield strength and ultimate compressive strength but also maintain a fracture strain above 19% with a 10 wt.% TiB₂@Ti addition, demonstrating more balanced and comprehensive mechanical properties. This strategy of optimizing microstructures through post-treatment processes rather than simply increasing the content of reinforcing phases provides a new path for the synergistic optimization of strength and toughness in HEAMCs, particularly in demanding applications such as aerospace and marine corrosion-resistant structural components requiring high comprehensive performance.

3.3.3. Strengthening Mechanism of Mechanical Property

The enhancement in mechanical properties of TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs induced by heat treatment originates from the synergistic effect of multi-scale strengthening mechanisms, with the core lying in the microstructure optimization and interface performance regulation driven by element diffusion. Combined with XRD, SEM/EDS, and mechanical test results, the strengthening mechanisms can be deconstructed into the following four levels:

(1)

Matrix strength improvement is dominated by Ti solid solution strengthening. XRD analysis shows that as the heat treatment temperature increases from 800 °C to 1000 °C, the diffraction angle of the (111) crystal plane of the FCC phase shifts leftward from 44.0° to 43.8°, corresponding to the lattice constant increasing from 0.3569 nm to 0.358 Å (Figure 1b). This change is attributed to the diffusion of Ti atoms (radius 1.47 Å) in TiB₂@Ti at the grain boundaries into the FCC matrix (Table 2 shows that the Ti content in the solid solution zone increases from 18.2 at.% to 25.1 at.%), and their larger atomic radius causes approximately 1.12% distortion of the matrix lattice, forming a stress field that hinders dislocation movement. According to the Gurao model, there is a linear relationship between the shear modulus increment ΔG caused by lattice distortion and the strength improvement Δσ:Δσ ≈ 2GΔε, where G is the shear modulus and Δε is the distortion amount. Combined with the shear modulus of the FCC phase G ≈ 76 GPa, it is calculated that the Ti solid solution contributes about 40% of the yield strength improvement (from 955 MPa to 1294 MPa), becoming the main strengthening mechanism.

(2)

Synergistic strengthening of grain boundary segregation and particle dispersion. EDS analysis indicates that heat treatment promotes the reverse diffusion of elements such as Ni and Co to the grain boundaries (the Ni content in the solid solution zone decreases from 31.2 at.% to 27.8 at.%), forming a solute segregation layer with a thickness of about 50–100 nm (Figure 2d). This segregation layer improves the grain boundary strength in two ways: 1) reducing the grain boundary energy (estimated to decrease by 15–20%) to inhibit crack initiation; and 2) forming a short-range ordered structure to increase the resistance of dislocations crossing the grain boundary. The grain boundary hardness at 1000 °C reaches 618.4 HV, increasing by 32.5% compared with the as-sintered state, corresponding to about 30% of the strengthening contribution. Atomic diffusion at high temperatures promotes the dispersion of TiB₂ particles from the agglomerates in the as-sintered state (Figure 2a) into particles with an average size of 2–3 μm (Figure 2d), and the particle spacing decreases from 5 μm to 3 μm. According to the Orowan mechanism, the critical shear stress required for dislocations to bypass the particles increases by about 20% at 1000 °C, corresponding to a strength contribution of about 20%.

(3)

Synergistic regulation of strength and plasticity by eutectic structure evolution. SEM observation shows that heat treatment transforms the eutectic zone from lamellar (Figure 2a) to island-like (Figure 2d), reducing the phase boundary area by about 30% and effectively decreasing the interface stress concentration. The island-like BCC phase (rich in Cr/Al/Ni/B) and the FCC matrix (rich in Co/Fe/Ni) form a “hard–soft” composite structure: the island-like phase acts as a strengthening phase to hinder dislocation slip (hardness ≈ 580 HV), while the FCC matrix provides a plastic deformation channel (elongation ≈ 19.4%). This structure makes the material exhibit the synergistic effect of “work hardening–crack deflection” during compression, and the fracture strength increases from 1910 MPa to 2385 MPa. Based on the binary mixing enthalpies (Ti-Ni −35 kJ·mol⁻¹, B-Cr −31 kJ·mol⁻¹ [29]), heat treatment induces the directional diffusion of elements: Ti is the preferential solid dissolved in the FCC phase to reduce the system energy, and B is enriched in the BCC phase to form stable compounds. This distribution avoids the interface weakening caused by the agglomeration of TiB₂ in the as-sintered state, enabling the composite to maintain plasticity at high strength.

(4)

Multi-mechanism synergy, breaking through the traditional material design bottleneck. Compared with the single-particle-reinforced TiB₂/AlCoCrFeNi_2.1 composite reported by Han et al. [13] (strength 2.5 GPa but elongation < 10%), this study achieves the simultaneous improvement of strength and plasticity (yield strength +35.5%, elongation 19.4%) through the fourfold mechanisms of “solid solution-grain boundary-particle-structure” activated by heat treatment.

4. Conclusions

This study investigates the effects of different heat treatment temperatures on the microstructure and mechanical properties of 10 wt.% TiB₂@Ti/AlCoCrFeNi_2.1 EHEAMCs. The main findings are as follows:

(1)

Both as-sintered and heat-treated composites consist of FCC, BCC, TiB₂, and Ti phases, with a stable preferred orientation of the (111) crystal plane in the FCC phase. As the temperature increases, the BCC phase diffraction peak separates from the main FCC peak and increases in intensity, while the peak positions of the FCC/BCC phases shift leftward, which is attributed to the diffusion of TiB₂@Ti from the grain boundaries into the matrix. The solid solution of Ti increases the lattice constant of the FCC phase. Microstructurally, the eutectic region transforms from lamellar to island-like structures, and the solid solution zone narrows. The black blocky substances at the grain boundaries are rich in Ti and B (Ti concentration > 64 at.%), while the Ti content in the solid solution zone increases with temperature (from 18.2 at.% to 25.1 at.%), and elements such as Ni decrease. Element diffusion is driven by binary mixing enthalpy, with Ti and B tending to solidify in the FCC and BCC phase regions, respectively.

(2)

Mechanical properties significantly improve with increasing temperature. The average hardness increases from 467.3 HV in the as-sintered state to 579.2 HV (+23.9%) at 1000 °C, with the grain boundary hardness exceeding 600 HV. Yield strength increases from 955 MPa to 1294 MPa (+35.5%), ultimate strength increases from 1910 MPa to 2385 MPa (+24.9%), and fracture strain remains above 19%. The strengthening mechanisms include lattice distortion induced by the Ti solid solution (solid solution strengthening), enhanced BCC phase strength via B-Cr binding, reduced grain boundary energy due to Ni/Co diffusion (grain boundary strengthening), improved efficiency of TiB₂ particle dispersion via the Orowan mechanism, and stress concentration reduction from eutectic islanding, collectively achieving synergistic enhancement of strength and plasticity.