Composite ceramic particle reinforced high-entropy alloy matrix composite and preparation method thereof

Abstract

This application discloses a composite ceramic particle-reinforced high-entropy alloy matrix composite material and its preparation method, belonging to the field of high-entropy alloy composite materials. The composite material includes a high-entropy alloy matrix and composite ceramic particles dispersed within the high-entropy alloy matrix. A nanocomposite interface layer exists between the composite ceramic particles and the high-entropy alloy matrix, which is generated in situ by bidirectional interdiffusion of elements between the high-entropy alloy matrix and the composite ceramic particles. The composite ceramic particles include TiC, TiN, ZrB2, TiB2, BN, and NbN, and the high-entropy alloy matrix includes any one of the following multi-element alloys: NiCoCrWMoFe, NiCoCrWMoAl, and NiCoCrWFeMnSi. The composite ceramic particle-reinforced high-entropy alloy matrix composite material of this application significantly improves high-temperature tensile strength and high-temperature creep resistance, making it suitable for manufacturing high-temperature hot-end components such as turbine blades for aero-engines.

Description

Technical Field

[0001] This application belongs to the field of high-entropy alloy composite materials, and specifically relates to a composite ceramic particle reinforced high-entropy alloy matrix composite material and its preparation method. Background Technology

[0002] Aero-engine turbine blades, combustion chambers, and other hot-end components operate under extreme high-temperature and high-stress environments for extended periods, placing near-stringent demands on the high-temperature strength and durability of materials. Employing ceramic-phase-reinforced metal matrices is a mainstream approach to improving component performance; however, existing technologies have significant limitations in addressing the core issue of interfacial bonding reliability, severely restricting their practical application potential.

[0003] First, while traditional single-ceramic reinforcing phase (such as TiC or WC) systems can improve the hardness and strength of the matrix, the interfacial bonding is weak due to the mismatch between the reinforcing phase and the metal matrix (such as nickel-based, titanium-based, and other high-temperature alloys). Under thermal stress and mechanical loads, this interface is extremely prone to becoming a weak link for crack initiation and propagation, leading to premature material failure under extreme conditions.

[0004] Secondly, some composite ceramic reinforcing phases used in existing technologies (such as Al2O3, CrC, and WC) also have inherent defects. Although these materials have high hardness, their physical compatibility with the high-temperature alloy matrix is ​​poor, and the mismatch in thermal expansion coefficients leads to high residual stress at the interface. More importantly, they usually only form a simple mechanical bond with the matrix, or undergo harmful interfacial reactions to generate brittle phases (such as certain oxides or carbides), making it difficult to form a strong, tough, and stable gradient transition interface layer in situ. This fragile bonding state will exacerbate interface degradation under high-temperature, long-term service, causing reinforcing phase delamination, thus failing to meet the stringent requirements of aero-engines for the long-term reliability of components.

[0005] In addition, the following problems are prone to occur in the existing traditional ceramic-reinforced metal matrix material preparation process: First, irregularly shaped ceramic particles are prone to stress concentration in the matrix, becoming crack sources; Second, the chemical compatibility between the ceramic phase and the metal matrix material is poor at high temperatures, the interface bond is fragile, and adverse interface reactions or dissolution are prone to occur during the rapid melting and solidification process of laser additive manufacturing, leading to a decrease in the reinforcement effect; Third, the uniformity of the distribution of the reinforcing phase in the matrix is ​​difficult to control, especially when using premixed powder for additive manufacturing, the difference in powder flowability and bulk density is prone to component segregation, which directly affects the density and performance uniformity of the formed parts. Summary of the Invention

[0006] To address the aforementioned problems, the main objective of this application is to provide a composite ceramic particle-reinforced high-entropy alloy matrix composite material and its preparation method. The material provided in this application possesses excellent high-temperature tensile strength and high-temperature creep resistance, making it suitable for manufacturing high-temperature hot-end components such as turbine blades for aero-engines. The method in this application achieves uniform dispersion and strong interfacial bonding of the ceramic phase in the matrix through spheroidized composite design and laser in-situ interface fabrication technology.

[0007] To achieve the above objectives, this application provides a composite ceramic particle-reinforced high-entropy alloy matrix composite material, comprising a high-entropy alloy matrix and composite ceramic particles dispersed in the high-entropy alloy matrix; the particle size of the composite ceramic particles is multi-scale, ranging from nanometer to micrometer.

[0008] The composite ceramic particles and the high-entropy alloy matrix have a nanocomposite interface layer, which is generated in situ by bidirectional interdiffusion of elements between the high-entropy alloy matrix and the composite ceramic particles.

[0009] The composite ceramic particles include TiC, TiN, ZrB2, TiB2, BN, and NbN, and the high-entropy alloy matrix includes any one of the following multi-element alloys: NiCoCrWMoFe, NiCoCrWMoAl, and NiCoCrWFeMnSi.

[0010] Furthermore, the composite ceramic particles account for 5-30% of the mass fraction of the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

[0011] Furthermore, the composite ceramic particles comprise the following components by mass percentage: TiC: 28%-30%, TiN: 30%-35%, BN: 10%-13%, ZrB2: 7%-8%, NbN: 3%-5%, TiB2: 9%-17%.

[0012] Furthermore, the loose packing density of the composite ceramic particles is 0.6-4.2 g / cm³. 3 The flow rate is 10-40 s / 50 g.

[0013] Furthermore, the thickness of the nanocomposite interface layer is 40-300 nm.

[0014] A second aspect of this application also provides a method for preparing the aforementioned composite ceramic particle-reinforced high-entropy alloy matrix composite material, comprising the following steps:

[0015] Step S1: Mix various ceramic materials such as TiC, TiN, ZrB2, TiB2, BN, and NbN with water and binder to form a slurry;

[0016] Step S2: The slurry is subjected to spray granulation, high-temperature sintering, and plasma spheroidization in sequence, and then screened to obtain the composite ceramic particles.

[0017] Step S3: The composite ceramic particles are mixed with high-entropy alloy matrix powder and then subjected to laser additive manufacturing to obtain the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

[0018] Furthermore, the laser additive manufacturing employs a laser selective melting process or a laser melting deposition process.

[0019] Furthermore, the parameters of the laser melting deposition process are set as follows: laser power 400-1500W, scanning speed 400-1200mm / min, powder feeding tray rotation speed 800-2800r / min, carrier gas flow rate 3-10L / min, and overlap rate 40-50%;

[0020] Furthermore, the parameters of the laser selective melting process are set as follows: laser power 200-350W, scanning speed 800-1500mm / s, scanning spacing 0.1-0.13mm, auxiliary powder layer thickness 30-50μm, and overlap rate 40-50%.

[0021] Furthermore, in step S3, after the laser additive manufacturing, the high-entropy alloy-based composite material after laser additive manufacturing is subjected to solution treatment and aging treatment in sequence under vacuum conditions.

[0022] Furthermore, in the solution treatment, the high-entropy alloy-based composite material is heated to 1250±10℃, held at that temperature for 2±0.25h, and then cooled to room temperature under an inert atmosphere.

[0023] Furthermore, in the aging treatment, the high-entropy alloy-based composite material after solution treatment is heated to 750±20℃, held at that temperature for 1±0.25h, and then cooled to room temperature by filling with an inert atmosphere.

[0024] Further, in step S2, the spray granulation includes: atomizing the slurry under a temperature field of 250-300°C inlet temperature and 100-105°C outlet temperature, and then drying the atomized droplets to form agglomerates;

[0025] Furthermore, the high-temperature sintering treatment includes: holding the spray-granulated agglomerates at 750-850℃ for 5-8 hours under a protective atmosphere, and then holding them at 1000-1400℃ for 6-8 hours to obtain the sintered product.

[0026] Furthermore, the parameter settings for the plasma spheroidization treatment include: controlling the plasma beam current to 620-630A, the main gas flow rate to 120-130 SCFH, the auxiliary gas flow rate to 10-11 SCFH, the carrier gas flow rate to 9-12 SCFH, and the powder feeding rate to 4.0-6.0 r / min.

[0027] Furthermore, the high-entropy alloy matrix powder is prepared by a method including the following steps: weighing each elemental metal raw material according to the target composition of the high-entropy alloy matrix, and performing vacuum induction melting on each weighed elemental metal raw material under a protective atmosphere to obtain an alloy melt;

[0028] The alloy melt is atomized under vacuum conditions to obtain spherical or near-spherical high-entropy alloy powder using high-pressure inert gas as the atomization medium.

[0029] The high-entropy alloy powder is sieved to obtain high-entropy alloy powder of the target particle size as high-entropy alloy matrix powder.

[0030] Compared with the prior art, this application has the following beneficial effects:

[0031] To address the performance bottleneck of hot-section components in aero-engines under extreme high temperatures, this application provides a comprehensive, chain-wide collaborative technical solution encompassing material structure design, composite ceramic particle preparation, and laser melting forming. By combining composite ceramic particles containing multiple ceramic phases with a high-entropy alloy metal matrix and implementing interfacial reactions, a stable microstructure with multi-level strengthening characteristics is constructed. Compared to existing technologies employing traditional reinforcing phases such as Al2O3, CrC, and WC, this solution achieves a systematic breakthrough in process adaptability, interfacial bonding quality, and high-temperature stability, specifically manifested in:

[0032] First, the composite ceramic particle system of this application exhibits unparalleled superiority over traditional methods in terms of high-temperature mechanical properties and long-term structural stability. The synergistic interaction among the selected multi-element ceramic phases (TiC, TiN, ZrB2, TiB2, BN, and NbN) not only provides high hardness and high wear resistance but also effectively enhances the high-temperature stability and self-lubricating properties of the material. Together, they construct a strong and tough composite reinforcement skeleton, which is dispersed in the high-entropy alloy matrix at multiple scales, generating a strong dispersion strengthening and pinning effect. Furthermore, it forms a thermodynamically stable reinforcement network with the high-entropy alloy matrix through an in-situ reaction-formed nanocomposite interface layer. This unique microstructure enables the composite material to effectively resist creep deformation and suppress interfacial diffusion and instability when subjected to high temperatures (e.g., above 1000°C) and high stress, thereby elevating high-temperature tensile strength and service life to levels that are difficult to achieve with traditional reinforced composite materials.

[0033] Secondly, regarding the interface bonding mechanism and strengthening effect, the in-situ interface reaction between the multi-component ceramic phase and the high-entropy alloy matrix in this application creates a strong and tough interface that traditional reinforcing phases cannot achieve. This application utilizes "laser in-situ interface technology," where a pre-fabricated multi-component composite ceramic phase undergoes controllable in-situ melting and precipitation under the action of a high-energy laser. This process not only forms a stable reinforcing phase with multi-scale dispersed distribution but also grows a nanocomposite interface layer at the interface, primarily composed of (Ti,W,Nb)C solid solution and Ti(C,N). This interface layer is firmly bonded, with a continuous transition in composition, and can efficiently transfer loads and inhibit crack propagation, fundamentally solving the problem of early failure caused by the fragile or brittle reaction layer of traditional reinforcing phases.

[0034] Third, regarding process adaptability and forming quality, the "spheroidized composite design" of this application fundamentally overcomes the problem of incompatibility between traditional reinforcing phases and additive manufacturing processes. This application prepares a specialized reinforcing phase with uniform composition and high sphericity by synergistically spheroidizing and composite-designing multi-element ceramic materials such as TiC, TiN, ZrB2, TiB2, BN, and NbN through spray granulation, high-temperature sintering, and plasma spheroidization. Its physical properties (density, flowability) perfectly match the high-entropy alloy matrix. This ensures excellent flow and spreadability of the composite powder, a prerequisite for achieving high-density, defect-free forming in laser additive manufacturing.

[0035] Fourth, the dual heat treatment system of "solution treatment + aging treatment" introduced after laser additive manufacturing further optimizes the microstructure of the material and fully taps its performance potential. Specifically, the solution treatment at around 1250℃ aims to fully dissolve and homogenize alloying elements (especially diffusion elements from the ceramic phase) in the high-entropy alloy matrix, eliminating internal stress and compositional segregation caused by rapid solidification, and laying a uniform supersaturated solid solution foundation for subsequent precipitation strengthening. The subsequent aging treatment at around 750℃ promotes the orderly and uniform precipitation of fine nanoscale second phases (which may include newly formed phases from matrix or interface reactions) in the supersaturated solid solution. These dispersed nanoscale precipitates can exert a strong pinning effect on dislocation movement, producing a significant precipitation strengthening effect, thereby enabling the composite material to achieve another qualitative leap in strength, hardness, and high-temperature stability without significantly impairing plasticity.

[0036] In summary, this application, through innovation in the material system and combined with the manufacturing processes of "spheroidized composite design", "laser in-situ interface technology" and "solution treatment + aging treatment", transforms the theoretical advantages of multi-component composite ceramic phases into engineerable, high-performance next-generation aero-engine hot-end components, achieving a breakthrough improvement in high-temperature mechanical properties.

[0037] Other features and effects of this application will be set forth in the following description, and in part will be obvious from the description or may be learned by practicing this application. Attached Figure Description

[0038] To more clearly illustrate the technical solutions in the embodiments of this application or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are some embodiments of this application. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0039] Figure 1 The microstructure diagram of the high-entropy alloy spherical powder of Embodiment 1 of this application is shown.

[0040] Figure 2 The microstructure diagram of the composite ceramic particles after plasma spheroidization treatment in Example 1 of this application is shown.

[0041] Figure 3 The microstructure diagrams of the composite ceramic particle-reinforced high-entropy alloy matrix composite material at different magnifications of Embodiment 1 of this application are shown; wherein, Figure 3 (a) is a scanning electron microscope image of the composite ceramic particle-reinforced high-entropy alloy matrix composite material at 2000 magnification. Figure 3 (b) is a scanning electron microscope image of the composite ceramic particle-reinforced high-entropy alloy matrix composite material at 10,000 magnification.

[0042] Figure 4 Image (a) shows a partial microstructure of the composite ceramic particle-reinforced high-entropy alloy matrix composite material of Example 1 of this application, as shown in a dark-field transmission electron microscope image. Figure 4 (b) in the middle is Figure 4 (a) Electron diffraction pattern at the selected location.

[0043] Figure 5 Image (a) shows another local microstructure of the composite ceramic particle-reinforced high-entropy alloy matrix composite material of Example 1 of this application, as shown in transmission electron microscopy dark-field photograph. Figure 5 (b) in the middle is Figure 5 (a) Electron diffraction pattern at the selected location. Detailed Implementation

[0044] To make the objectives, technical solutions, and advantages of the embodiments of this application clearer, the technical solutions in the embodiments of this application will be clearly and completely described below in conjunction with the embodiments of this application. Obviously, the described embodiments are only some embodiments of this application, not all embodiments. Based on the embodiments in this application, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of this application.

[0045] To achieve the above objectives, a first aspect of the embodiments of this application provides a composite ceramic particle-reinforced high-entropy alloy matrix composite material, comprising a high-entropy alloy matrix and composite ceramic particles dispersed in the high-entropy alloy matrix; the particle size of the composite ceramic particles is multi-scale, ranging from nanometer to micrometer.

[0046] The composite ceramic particles and the high-entropy alloy matrix have a nanocomposite interface layer, which is generated in situ by bidirectional interdiffusion of elements between the high-entropy alloy matrix and the composite ceramic particles.

[0047] The composite ceramic particles include TiC, TiN, ZrB2, TiB2, BN, and NbN, and the high-entropy alloy matrix includes any one of the following multi-element alloys: NiCoCrWMoFe, NiCoCrWMoAl, and NiCoCrWFeMnSi.

[0048] To fundamentally solve the aforementioned interface problems, this application creatively constructs composite ceramic particles with multi-scale (nanometer to micrometer scale, approximately 30 nm-3 μm in some optional embodiments) dispersed distribution within a high-entropy alloy matrix. These composite ceramic particles form multiple stable ceramic phases within the matrix, generating strong dispersion strengthening and pinning effects, significantly enhancing the interfacial stability between the ceramic phase and the matrix. Specifically, the composite ceramic particles of this application employ a multi-component composite ceramic phase comprising TiC, TiN, ZrB2, TiB2, BN, and NbN as the reinforcing body. This system integrates the high hardness and wear resistance of TiC / TiN, the high-temperature strength and thermal shock resistance of TiB2, the high melting point of NbN, the self-lubricating properties and stability of BN, and the introduced ZrB2 as a core component of ultra-high temperature ceramics. Its extremely high melting point, excellent high-temperature stability, high hardness, and high modulus further solidify the load-bearing capacity of the reinforcing skeleton under extreme environments. In summary, these ceramic phases work synergistically and are indispensable to each other, jointly constructing a strong and tough composite material reinforcement skeleton, which exhibits superior performance in high-temperature mechanical properties and long-term structural stability compared to traditional methods.

[0049] Furthermore, controlled interdiffusion and interfacial reaction occurred at the interface between the composite ceramic particles and the high-entropy alloy matrix. Elements such as W and Al in the matrix diffused bidirectionally with elements such as Ti, C, N, and Zr in the ceramic particles, forming a nanocomposite interfacial layer in situ. This nanocomposite interfacial layer can effectively alleviate thermal mismatch stress caused by the difference in thermal expansion coefficients, inhibit the initiation and propagation of cracks at the interface, and effectively transfer external loads from the high-entropy alloy matrix to the ceramic reinforcing phase, thus achieving a simultaneous leapfrog improvement in the strength, toughness, and high-temperature durability of the composite material.

[0050] In summary, this application effectively overcomes the interfacial bonding problem that has long constrained the development of hot-end component materials through the synergistic design of multi-element ceramic phases and the successful construction of nanocomposite interface layers, providing an ideal material solution for the development of a new generation of high-performance aero-engines.

[0051] In a preferred embodiment of this application, the composite ceramic particles account for 5-30% of the mass fraction of the composite ceramic particle-reinforced high-entropy alloy matrix composite material. Extensive experimental verification has shown that this range of composite ceramic particle proportions allows the ceramic reinforcing phase to form an effective dispersed strengthening network in the matrix without significantly sacrificing matrix toughness and forming process stability. More preferably, the composite ceramic particles account for 15-25% of the mass fraction of the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

[0052] In a preferred embodiment of this application, the composite ceramic particles comprise the following components by mass percentage: TiC: 28%-30%, TiN: 30%-35%, BN: 10%-13%, ZrB2: 7%-8%, NbN: 3%-5%, and TiB2: 9%-17%. By specifically defining the specific component ratios (TiC, TiN, ZrB2, TiB2, BN, NbN) of the composite ceramic particles, the complementary advantages of multiple ceramic phases are achieved. This multi-component design not only provides high hardness and high wear resistance but also utilizes phases such as BN and ZrB2 to enhance the high-temperature stability and self-lubricating properties of the material, thereby jointly improving the overall high-temperature performance of the composite material.

[0053] In some preferred embodiments of this application, the loose packing density of the composite ceramic particles is 0.6-4.2 g / cm³. 3 Preferably, it is 1.0-1.8 g / cm³. 3The flow rate is 10-40 s / 50 g, preferably 20-40 s / 50 g. By specifically defining the density and flowability of the composite ceramic particles, a perfect match with the high-entropy alloy matrix can be ensured, guaranteeing excellent uniformity and flowability in the subsequent mixing process. This results in highly dense, defect-free products in the subsequent laser additive manufacturing process, solving the technical problems of segregation and poor flowability of traditional irregular ceramic particles in additive manufacturing. Furthermore, the particle size of the high-entropy alloy matrix is ​​15-120 μm.

[0054] In a preferred embodiment of this application, the thickness of the nanocomposite interface layer is 40-300 nm. Further, the composition of the nanocomposite interface layer exhibits a gradient-continuous change: the concentration of elements from the high-entropy alloy matrix gradually decreases from the alloy matrix side to the ceramic particle side, while the concentration of elements from the ceramic particles gradually increases from the alloy matrix side to the ceramic particle side.

[0055] The nanocomposite interface layer of this application comprises a multi-component composite compound formed by the interdiffusion reaction of high-entropy alloy matrix elements and ceramic phases (such as TiC, TiN, ZrB2, NbN, BN, TiB2). This multi-component composite compound forms a series of complex solid solutions, including (Ti, Al, Zr, Nb)(C, N), Cr(C, N), (Ti, Cr)(C, N), (Ti, Al)(C, N), and Ti(C, N), among others, with gradually varying compositions. Notably, boron (B) is primarily dissolved in stable borides or composite phases, without forming an independent B-rich interface phase.

[0056] The composite ceramic particle-reinforced high-entropy alloy matrix composite material of this application possesses a unique multiphase microstructure: on the one hand, a strong and tough metallurgical bond is achieved between the ceramic particles and the matrix through the aforementioned gradient nanocomposite interface layer; on the other hand, some incompletely dissolved or dissolved-precipitated original ceramic particles (such as TiC, TiN, ZrB2, NbN, BN, TiB2) are still retained in the microstructure. This synergistic effect of the gradient interface layer and the residual ceramic particles together constitutes the basis for the excellent comprehensive performance of the composite material.

[0057] A second aspect of the embodiments of this application also provides a method for preparing the aforementioned composite ceramic particle-reinforced high-entropy alloy matrix composite material, comprising the following steps:

[0058] Step S1: Mix various ceramic materials such as TiC, TiN, ZrB2, TiB2, BN, and NbN with water and binder to form a slurry;

[0059] Step S2: The slurry is subjected to spray granulation, high-temperature sintering, and plasma spheroidization in sequence, and then screened to obtain the composite ceramic particles.

[0060] Step S3: The composite ceramic particles are mixed with high-entropy alloy matrix powder and then subjected to laser additive manufacturing to obtain the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

[0061] The composite ceramic particle-reinforced high-entropy alloy matrix composite material of this application is not simply a matter of combining ceramic particles with a high-entropy alloy matrix. First, various carbon / nitrogen / borides (TiC, TiN, TiB2, BN, NbN, ZrB2, etc.) are prepared into a slurry, which is then subjected to spray granulation, high-temperature sintering, and plasma spheroidization. The resulting composite ceramic particles exhibit high sphericity, high density, and good flowability, and their physical properties match those of the high-entropy alloy matrix. In contrast, traditional ceramic phase powders such as Al2O3, CrC, and WC often have irregular shapes and poor flowability, easily leading to compositional segregation, porosity, and incomplete fusion defects during laser processing, severely limiting the reliability of additive manufacturing processes in complex and precision components. Furthermore, Al2O3 may react adversely with active elements in the matrix at high temperatures, and WC and CrC are prone to oxidation and grain growth during long-term high-temperature service, resulting in performance degradation.

[0062] The multi-component ceramic phases (TiC, TiN, ZrB2, TiB2, BN, NbN) selected in this application possess excellent high-temperature stability. Under the action of a high-energy beam processed by laser melting and forming, the pre-formed spheroidized composite ceramic particles do not undergo mechanical mixing, but rather controlled in-situ melting and precipitation, forming a strong and tough interface that traditional reinforcing phases cannot achieve. Traditional reinforcing phase Al2O3 is chemically stable but has poor wettability with most metal matrices, resulting in interfaces that are mostly mechanically bonded with weak bonding forces. While WC and CrC can react with some metals, they often form continuous, brittle interfacial compounds (such as W2C or complex carbides), which are prone to becoming crack initiation sites under thermal stress.

[0063] Under the irradiation of high-energy lasers, the multi-component composite ceramic particles of this application not only form a stable reinforcing phase with a multi-scale dispersed distribution, but also grow in situ a nanocomposite interface layer mainly composed of (Ti,W,Nb)C solid solution and Ti(C,N). This interface layer is firmly bonded and has a continuous transition in composition, enabling the composite material to effectively resist creep deformation and suppress interface diffusion and instability when subjected to high temperatures (such as above 1000°C) and high stress. This solves the problem of early failure caused by the fragile interface or brittle reaction layer of traditional reinforcing phases, thereby improving the high-temperature tensile strength and service life to a level that is difficult for traditional reinforced composite materials to achieve.

[0064] In a preferred embodiment of this application, in step S1, the amount of water added is 1.5-1.8 times the total mass of the ceramic raw materials, and the amount of binder added is 0.15-0.3% of the total mass of the ceramic raw materials.

[0065] In a preferred embodiment of this application, step S2, the spray granulation includes: atomizing the slurry at a rotation speed of 45-50 rpm under a temperature field of 250-300°C inlet temperature and 100-105°C outlet temperature, while simultaneously drying the droplets to form agglomerates at a feed rate of 20-25 rpm. Under these conditions, the slurry is atomized into microdroplets and rapidly dried during its descent, forming porous agglomerated powder with uniform composition and approximately spherical shape, providing a qualified intermediate product for subsequent sintering and spheroidization treatments.

[0066] In a preferred embodiment of this application, step S2, the high-temperature sintering treatment includes: holding the spray-granulated agglomerates at 750-850℃ for 5-8 hours under a protective atmosphere to fully remove the binder, followed by holding at 1000-1400℃ for 6-8 hours to allow diffusion bonding between ceramic particles, achieving preliminary densification and obtaining a sintered product. To facilitate subsequent plasma spheroidization treatment, the sintered product is then sieved for preliminary particle size control. The sieving is performed using a 240-600 mesh sieve, with a single processing volume of 500-800g, an amplitude of 3-5mm, and a time of 10-20 minutes.

[0067] In a preferred embodiment of this application, to further improve sphericity, density, and surface finish, the sintered intermediate particles are fed into a plasma spheroidizing device for final processing. The parameters for plasma spheroidizing are: plasma beam current of 620-630 A, main gas flow rate of 120-130 SCFH, auxiliary gas flow rate of 10-11 SCFH, carrier gas flow rate of 9-12 SCFH, and powder feeding rate of 4.0-6.0 r / min. The particles melt instantaneously in the high-temperature jet of the plasma beam, forming perfectly spherical droplets under surface tension, and are then rapidly cooled and solidified. The spheroidized powder is then precision sieved again using a 270-500 mesh sieve. Under conditions of 200-300 g per batch, amplitude of 5-7 mm, and time of 30-40 minutes, a final loose density of 0.6-4.2 g / cm³ is obtained. 3 High-quality spheroidized composite ceramic particles with a flowability of 10-40s / 50g.

[0068] In some preferred embodiments of this application, the laser additive manufacturing can be carried out using selective laser melting (SLM) or laser melting deposition (LMD) additive manufacturing methods.

[0069] In the laser melting deposition process, plasma-spheroidized composite ceramic particles and high-entropy alloy matrix powder are simultaneously transported through a dual-channel powder feeding system, or the composite powder formed by the plasma-spheroidized composite ceramic particles and high-entropy alloy matrix powder is transported through a single-channel powder feeding system (premixing + single-channel powder feeding) for laser melting deposition. The parameters of the laser melting deposition process are set as follows: laser power 400-1500W, scanning speed 400-1200mm / min, powder feeding tray rotation speed 800-2800r / min, carrier gas flow rate 3-10L / min, and overlap rate 40-50%.

[0070] In the selective laser melting (SLM) process, plasma-spheroidized composite ceramic particles are pre-mixed with high-entropy alloy matrix powder to form a composite powder. This composite powder is then formed layer by layer through selective laser melting using a powder-layout SLM device. The process parameters are set as follows: laser power 200-350W, scanning speed 800-1500mm / s, scanning spacing 0.1-0.13mm, auxiliary powder layer thickness 30-50μm, and overlap rate 40-50%. Under high-purity argon protection, the laser beam selectively melts each layer of powder according to the path of the three-dimensional model, accumulating layer by layer to form a dense component.

[0071] In a preferred embodiment of this application, step S3, after the laser additive manufacturing, further includes: subjecting the laser-added high-entropy alloy-based composite material to solution treatment and aging treatment sequentially under vacuum conditions. Exemplarily, in the solution treatment, the high-entropy alloy-based composite material is heated to 1250±10℃, held at that temperature for 2±0.25h, and then cooled to room temperature under an inert atmosphere. Exemplarily, in the aging treatment, the solution-treated high-entropy alloy-based composite material is heated to 750±20℃, held at that temperature for 1±0.25h, and then cooled to room temperature under an inert atmosphere.

[0072] For composite materials formed by laser additive manufacturing, their unique non-equilibrium rapid solidification microstructure is more sensitive to heat treatment. Traditional homogenization annealing may not effectively eliminate manufacturing defects (such as micropores and lack of fusion), and may instead lead to abnormal grain growth or deterioration of reinforcing phase distribution due to thermal cycling. Therefore, this application develops a heat treatment process matched with a specific ceramic-high-entropy alloy system to synergistically optimize the matrix microstructure, stabilize the interface structure, and release manufacturing internal stress. First, solution treatment eliminates the microsegregation and internal stress caused by rapid solidification during additive manufacturing, making the high-entropy alloy matrix composition more uniform and promoting more sufficient interdiffusion of components in the nanocomposite interface layer between the composite ceramic particles and the matrix, resulting in a stronger and more continuous interface bond. Second, the subsequent aging treatment induces the dispersion precipitation of a large number of nanoscale secondary carbonitrides (rich in elements such as Ti, Nb, and W) in the matrix. These precipitated phases, together with the original multi-scale composite ceramic particles, constitute a multi-level reinforcing phase architecture, which has an extremely strong synergistic inhibitory effect on dislocation movement. This specific heat treatment process significantly improves the overall high-temperature performance of composite ceramic particle-reinforced high-entropy alloy matrix composites.

[0073] In a preferred embodiment of this application, the high-entropy alloy matrix powder is prepared by a method comprising the following steps: weighing each elemental metal raw material according to the target composition of the high-entropy alloy matrix; performing vacuum induction melting on the weighed elemental metal raw materials under a protective atmosphere to obtain an alloy melt; atomizing the alloy melt under vacuum conditions using a high-pressure inert gas as the atomizing medium to obtain spherical or near-spherical high-entropy alloy powder; and sieving the high-entropy alloy powder to obtain high-entropy alloy powder of the target particle size as the high-entropy alloy matrix powder. Further, the particle size of the high-entropy alloy powder of the target particle size is 15-120 μm.

[0074] The fabrication process described in this application achieves the following results: Through precise coordination of process parameters throughout the entire fabrication process, the repeatable and stable manufacturing of high-performance microstructures is ensured. From the vacuum low-oxygen atomization of the matrix powder, the formulation, granulation, sintering, and spheroidization of the ceramic slurry, to the precise energy input for laser additive manufacturing, and the solution treatment and aging heat treatment after laser additive manufacturing, a complete and closed-loop process window is provided. These parameters support each other, ensuring that in-situ reactions, microstructure refinement, and interface strengthening occur as designed during the process from high-quality raw materials to the final component forming, ultimately obtaining composite material parts with uniformly dispersed ceramic phases, strong and tough interfaces, and virtually no defects.

[0075] The present application will be further described in detail below with reference to specific embodiments, which should not be construed as limiting the scope of protection claimed in the present application.

[0076] Example 1

[0077] A method for preparing a composite ceramic particle-reinforced high-entropy alloy matrix composite material includes the following steps:

[0078] (1) Preparation of high-entropy alloy matrix powder: NiCoCrWMoFe hexa-element high-entropy alloy powder was prepared by vacuum induction melting gas atomization process, including:

[0079] 1) Weigh the following elemental metal raw materials by weight percentage: nickel 22%, cobalt 29%, chromium 17%, tungsten 15%, molybdenum 9%, and iron 8%. Put the prepared raw materials into a vacuum induction melting furnace, evacuate to below 0.1 Pa, and then fill with high-purity argon as a protective atmosphere for induction heating and melting to obtain an alloy melt.

[0080] 2) Transfer the alloy melt to the atomization chamber, maintaining a vacuum level below 0.1 Pa. A stable liquid flow is formed through a guide tube, and high-pressure argon gas (5.5 MPa) is used as the atomization medium to impact and pulverize the melt into micron-sized droplets. These droplets rapidly cool and solidify in an inert atmosphere, forming spherical or near-spherical powder particles (see...). Figure 1 ).

[0081] 3) Finally, the powder particles are collected and sieved to obtain a high-entropy alloy matrix powder with a particle size of 15-53 μm, an oxygen content of 0.008%, and good sphericity.

[0082] (2) Preparation of composite ceramic particles:

[0083] 1) Raw material preparation and mixing: Weigh the ceramic raw materials according to the following mass ratios: TiC (290g, irregular particles with a particle size range of 1-3μm), TiN (350g, irregular particles with a particle size range of 2-4μm), BN (110g, irregular particles with a particle size range of 5-8μm), ZrB2 (75g, irregular particles with a particle size range of 3-10μm), NbN (40g, irregular particles with a particle size range of 2-6μm), and TiB2 (135g, irregular particles with a particle size range of 1-6μm). Mix the above ceramic raw materials with 1.6kg of deionized water and 1.8g of polyvinyl alcohol binder and ball mill to obtain a homogeneous slurry.

[0084] 2) Spray granulation: The slurry is conveyed to a spray drying tower for granulation. The inlet temperature is controlled at 280℃, the outlet temperature at 103℃, the spray disc rotation speed at 48 rpm, and the feed rate at 22 rpm. Under these conditions, the slurry is atomized into microdroplets and rapidly dried during its descent, forming porous aggregates with uniform composition and approximately spherical shape.

[0085] 3) High-temperature sintering treatment: Under argon protection, the spray-granulated agglomerates are first kept at 800℃ for 6 hours, and then kept at 1200℃ for 7 hours to obtain the sintered product.

[0086] 4) Plasma spheroidization treatment: After sintering, the sintered products are sieved and then subjected to plasma spheroidization treatment. The specific parameters are: current 625A, main gas (Ar) 125 SCFH, auxiliary gas (H2) 10.5 SCFH, carrier gas (Ar) 10 SCFH, and powder feeding rate 5.0r / min.

[0087] 5) After sieving, the particle size was 26-63 μm and the loose packing density was 1.7 g / cm³. 3 High-quality spherical composite ceramic particles with a flowability of 28 s / 50g (see [reference]). Figure 2 ).

[0088] (3) SLM laser selective melting process:

[0089] 1) The composite ceramic particles prepared above were dry-mixed with high-entropy alloy matrix powder at a mass ratio of 5:95 for 2 hours to obtain uniform composite powder.

[0090] 2) The above composite powder was subjected to selective laser melting (SLM) using an EOS M 250 device to form a composite ceramic particle-reinforced high-entropy alloy matrix composite material layer by layer. The process parameters were: laser power 280W, scanning speed 1200mm / s, scanning spacing 0.1mm, powder layer thickness 30μm, and protective atmosphere of high-purity argon (oxygen content <100 ppm).

[0091] Performance results:

[0092] The composite ceramic particle-reinforced high-entropy alloy matrix composite material prepared in this embodiment has a density of 99.5%, and the microstructure is shown by scanning electron microscopy. The composite ceramic phases (TiC, TiN, ZrB2, TiB2, BN, NbN) are dispersed in a multi-scale morphology of 50nm-2μm (see [reference]). Figure 3 In (a) and (b) of these examples, the thickness of the nanocomposite interface layer formed between the composite ceramic particles and the high-entropy alloy matrix is ​​approximately 80-150 nm. The main phases of this nanocomposite interface layer are (Ti,W,Nb)C solid solution and Ti(C,N) solid solution, etc. (see [reference]). Figure 4 and Figure 5 ), Figure 4 (a) is a dark-field transmission electron microscopy image of the microstructure of a low-magnification composite ceramic particle reinforced high-entropy alloy matrix composite material. It can be seen that a core-shell structure is formed in the microstructure. The core is mainly the ceramic particle reinforcing phase TiC, and the shell is an interface layer containing the reinforcing phase element. Figure 4 (b) is Figure 4 The diffraction pattern of the nanocomposite interface layer in selected area (a) shows that the interface layer in selected area is a (Ti,W,Nb)C solid solution phase. Figure 5(a) is a dark-field transmission electron microscopy image of the microstructure of the composite ceramic particle reinforced high-entropy alloy matrix composite material at high magnification. It can be seen that a core-shell structure is formed in the microstructure. The core is mainly the ceramic particle reinforcing phase TiN, and the shell is an interface layer containing the reinforcing phase element. Figure 5 (b) is Figure 5 The electron diffraction pattern of the nanocomposite interface layer in selected area (a) shows that the interface layer in this selected area is a Ti(C,N) solid solution phase.

[0093] The composite ceramic particle reinforced high-entropy alloy matrix composite material prepared above was subjected to high-temperature tensile strength test at 800℃ according to GB / T228.2-2015 "Metallic materials, tensile testing - Part 2: High temperature test method". Its high-temperature tensile strength reached 850MPa, which is about 37% higher than that of the single high-entropy alloy matrix without composite ceramic particles (620MPa).

[0094] Example 2

[0095] A method for preparing a composite ceramic particle-reinforced high-entropy alloy matrix composite material differs from Example 1 in that the amount of composite ceramic particles added is increased to 15 wt.%, and a premixed single-channel powder-feed laser melting deposition (LMD) process is used for laser melting and forming. This embodiment aims to achieve better high-temperature creep performance by increasing the reinforcing phase content. It includes the following steps:

[0096] (1) Preparation of high-entropy alloy matrix powder: The preparation process is the same as in Example 1, but the sieve particle size is adjusted to 53-106μm.

[0097] (2) Preparation of composite ceramic particles: The preparation process is the same as in Example 1, but the composition ratio is slightly adjusted to TiC (300g), TiN (350g), BN (100g), ZrB2 (80g), NbN (40g), and TiB2 (130g). The final sieved particle size is 26-63μm, and the loose packing density is 1.3 g / cm³. 3 High-quality spherical composite ceramic particles with a flowability of 30 s / 50g.

[0098] (3) Single-channel LMD laser melting deposition process:

[0099] 1) The composite ceramic particles prepared above are mixed with the high-entropy alloy matrix powder at a mass ratio of 15:85 to obtain composite powder.

[0100] 2) Composite ceramic particle-reinforced high-entropy alloy matrix composites were prepared by single-pass multilayer deposition (1000 layers) on a pre-fabricated high-entropy alloy substrate using an LMD (TruLaser Cell 3000) equipment with a single-channel powder feeding system. Process parameters: laser power 1000W, scanning speed 800 mm / min, powder feed tray speed 1500 r / min, powder carrier gas flow rate 6 L / min, overlap ratio (the proportion of the subsequent layer covering the previous layer) 40%.

[0101] Performance Results: The composite ceramic particle-reinforced high-entropy alloy matrix composite material deposited in this embodiment has a dense structure and a porosity of less than 0.5%. Microstructure analysis revealed that the composite ceramic phase is dispersed in a multi-scale morphology of 0.3-2.1 μm, forming a more continuous nanocomposite interface layer with a thickness of 150-250 nm. The main phases of the nanocomposite interface layer are (Ti,W)C solid solution and Ti(C,N) solid solution.

[0102] According to DL / T 2888-2025 "High-Temperature Creep Duration Test Procedure for Micro-Specimens of Metallic Materials", this embodiment achieved a fracture life of 100h under 750MPa stress in a high-temperature creep rupture test at 900℃, which far exceeds that of high-entropy alloy matrix materials (<30h), demonstrating the great potential of this material for long-term service at high temperatures.

[0103] Example 3

[0104] A method for preparing a ceramic-reinforced high-entropy alloy material differs from Example 1 in that the amount of composite ceramic particles added is increased to 25 wt%, and a dual-channel synchronous powder feeding laser melting deposition process is employed. This method aims to apply the material to extreme applications such as the leading edge of turbine blades to meet the extreme requirements of extremely high temperature strength and durability. The method includes the following steps:

[0105] (1) Preparation of high-entropy alloy matrix powder: The preparation process is the same as in Example 1, but the sieve particle size is adjusted to 53-106μm.

[0106] (2) Preparation of composite ceramic particles: The proportions of various ceramic raw materials were adjusted to TiC (280g), TiN (350g), BN (120g), ZrB2 (80g), NbN (50g), and TiB2 (120g). The power was increased during plasma spheroidization treatment, with the following parameters: current 630A, main gas (Ar) 130 SCFH, and the rest the same as in Example 1. The composite ceramic particles prepared in this example had higher sphericity, a flowability of 18s / 50g, a particle size of 26-63μm, and a loose packing density of 2.1g / cm³. 3 .

[0107] (3) Dual-channel LMD laser melting deposition process:

[0108] 1) The composite ceramic particles and high-entropy alloy matrix powder prepared above were fed into their respective independent powder feeders at a mass ratio of 25:75. Through a dual-channel powder feeding system, the composite ceramic particles and high-entropy alloy matrix powder were mixed and deposited in the laser molten pool of an LMD (TruLaser Cell 3000) device. The process parameters were: laser power 1300W, scanning speed 500 mm / min, powder feeder rotation speed 1800 r / min, powder-carrying gas flow rate 5 L / min, and overlap rate 45%.

[0109] Performance Results: The high-entropy alloy matrix composite material molded part reinforced with composite ceramic particles prepared in this embodiment maintained a high-temperature tensile strength of 550 MPa in a high-temperature tensile test at 1100℃ according to GB / T 228.2-2015 standard. In a high-temperature creep rupture test at 900℃ according to DL / T2888-2025 "Specification for High-Temperature Creep Testing of Micro-Specimens of Metallic Materials", its fracture life reached 180 hours under a stress of 750 MPa. Microstructural observation of the high-entropy alloy material molded part in this embodiment revealed that the extremely high ceramic content and dual-channel process produced richer nanoscale ceramic precipitates and a uniformly thick (approximately 200-300 nm) nanocomposite interface layer. This formed nanocomposite interface layer is the fundamental reason for achieving the extraordinary high-temperature performance and creep rupture properties.

[0110] Example 4

[0111] In this embodiment, the composite ceramic particle-reinforced high-entropy alloy matrix composite material obtained in Example 1 is used as the blank for subsequent heat treatment, specifically as follows:

[0112] The composite ceramic particle-reinforced high-entropy alloy matrix composite sample formed by SLM process in Example 1 was placed in a vacuum heat treatment furnace and post-processed according to the following steps: Solution treatment: After the vacuum level in the furnace was evacuated to below 0.1 Pa, high-purity argon was introduced as a protective atmosphere. The temperature was increased to 1250°C at a rate of 10°C / min and held at this temperature for 2 hours to fully homogenize the internal structure of the material and dissolve some of the non-equilibrium phases. After the holding period, high-purity argon was immediately introduced into the furnace for forced cooling to room temperature. Aging treatment: After the solution treatment, the furnace was evacuated to a low vacuum again and argon was introduced for protection. The temperature was increased to 750°C at a rate of 5°C / min and held for 1 hour to promote the precipitation of fine, dispersed secondary strengthening phases from the supersaturated solid solution. After the holding period, argon was introduced for cooling to room temperature.

[0113] Performance Results and Analysis:

[0114] Mechanical properties: After the above solution and aging treatments, the material's properties were significantly improved. A high-temperature tensile test was conducted at 800℃ according to GB / T 228.2-2015 standard, and the high-temperature tensile strength reached 920 MPa, an increase of approximately 8.2% compared to Example 1 without post-treatment (850 MPa), and an increase of approximately 48.4% compared to the pure matrix material (620 MPa). Simultaneously, the material's room-temperature toughness was also improved.

[0115] Microstructure Mechanism Analysis: Post-processing can play a crucial role in optimizing the microstructure of materials. First, solution treatment eliminates microsegregation and internal stress caused by rapid solidification of SLM, resulting in a more uniform composition of the high-entropy alloy matrix and promoting more complete interdiffusion of components in the nanocomposite interface layer (approximately 80-150 nm thick) between the composite ceramic particles and the matrix, leading to a stronger and more continuous interfacial bond. Second, the subsequent aging treatment induces the dispersion precipitation of a large number of nanoscale secondary carbonitrides (rich in elements such as Ti, Nb, and W) in the matrix. These precipitated phases, together with the original multi-scale composite ceramic particles, constitute a multi-level strengthening phase architecture, exerting an extremely strong synergistic inhibitory effect on dislocation movement.

[0116] Conclusion: This embodiment not only optimizes the matrix structure and interface structure through the above-mentioned specific post-processing process, but also introduces nanoscale secondary precipitation strengthening, thereby achieving a leap from "single-scale particle strengthening" to "multi-level synergistic strengthening", which significantly improves the overall high-temperature performance of the material, surpassing the improvement achieved by simply adjusting the amount of ceramic additives (such as in Examples 2 and 3).

[0117] Example 5

[0118] A method for preparing a composite ceramic particle-reinforced high-entropy alloy matrix composite material, differing from Example 1 in that the high-entropy alloy matrix is ​​replaced with a NiCoCrWMoAl multi-element alloy, specifically including the following steps:

[0119] (1) Preparation of high-entropy alloy matrix powder: NiCoCrWMoAl hexa-element high-entropy alloy powder was prepared by vacuum induction melting gas atomization process.

[0120] 1) Weigh the following elemental metal raw materials by weight percentage: nickel 22%, cobalt 29%, chromium 17%, tungsten 15%, molybdenum 9%, and aluminum 8%. Put the prepared raw materials into a vacuum induction melting furnace, evacuate to below 0.1 Pa, and then fill with high-purity argon as a protective atmosphere for induction heating and melting to obtain an alloy melt.

[0121] 2) The alloy melt was transferred to the atomization chamber and powder was prepared using the same high-pressure argon atomization process (5.5 MPa) as in Example 1.

[0122] 3) Finally, the powder particles are collected and sieved to obtain high-entropy alloy matrix powder with a particle size of 15-53μm, low oxygen content, and good sphericity.

[0123] (2) Preparation of composite ceramic particles: This step was performed entirely according to the method and proportions described in Example 1 (i.e., TiC: 29%, TiN: 35%, BN: 11%, ZrB2: 7.5%, NbN: 4%, TiB2: 13.5%), to prepare particles with a diameter of 26-63 μm and a bulk density of 1.7 g / cm³. 3 Spherical composite ceramic particles with a flowability of 28 s / 50g.

[0124] (3) SLM laser selective melting process:

[0125] 1) The composite ceramic particles prepared above were dry-mixed with NiCoCrWMoAl high-entropy alloy matrix powder at a mass ratio of 5:95 for 2 hours to obtain uniform composite powder.

[0126] 2) Using the same SLM process parameters as in Example 1 (laser power 280W, scanning speed 1200 mm / s, scanning spacing 0.1 mm, powder layer thickness 30μm), layer-by-layer forming was carried out to obtain composite ceramic particle reinforced high-entropy alloy matrix composite material.

[0127] Performance Results: The material prepared in this embodiment has high density, with well-dispersed composite ceramic particles in the microstructure, forming a nanocomposite interface layer with a continuously varying compositional gradient with the matrix. A high-temperature tensile test was conducted at 800℃ according to GB / T228.2-2015 standard, and its high-temperature tensile strength reached 920 MPa. This strength is significantly higher than that of the NiCoCrWMoAl-based composite material in Example 1. This is mainly attributed to the improved oxidation resistance brought about by the formation of a dense oxide film by aluminum (Al) at high temperatures, and the synergistic effect of molybdenum (Mo) with aluminum, chromium, and other elements further enhancing the high-temperature solid solution strength of the matrix.

[0128] Example 6

[0129] A method for preparing a composite ceramic particle-reinforced high-entropy alloy matrix composite material, differing from Example 1 in that the high-entropy alloy matrix is ​​replaced with a NiCoCrWFeMnSi multi-element alloy, specifically including the following steps:

[0130] (1) Preparation of high-entropy alloy matrix powder: NiCoCrWFeMnSi seven-element high-entropy alloy powder was prepared by vacuum induction melting gas atomization process.

[0131] 1) Weigh the following elemental metal raw materials by weight percentage: nickel 22%, cobalt 29%, chromium 17%, tungsten 15%, iron 8%, manganese 8%, and silicon 1%. Put the prepared raw materials into a vacuum induction melting furnace, evacuate to below 0.1 Pa, and then fill with high-purity argon as a protective atmosphere for induction heating and melting to obtain an alloy melt.

[0132] 2) The powder was prepared using the same high-pressure argon atomization process (5.5 MPa) as in Example 1.

[0133] 3) Collect and sieve the powder to obtain high-entropy alloy matrix powder with a particle size of 15-53μm.

[0134] (2) Preparation of composite ceramic particles: This step is carried out in accordance with the method and proportion described in Example 1 to prepare spherical composite ceramic particles of the same specifications.

[0135] (3) SLM laser selective melting process:

[0136] 1) Mix the composite ceramic particles with NiCoCrWFeMnSi high-entropy alloy matrix powder at a mass ratio of 5:95.

[0137] 2) The final material was obtained by layer-by-layer forming using the same SLM process parameters as in Example 1.

[0138] Performance results:

[0139] The material prepared in this embodiment has high density and uniform microstructure. The addition of manganese (Mn) helps stabilize the austenitic structure and improve the solid solution strengthening effect, while silicon (Si) enhances the material's oxidation resistance and casting fluidity. Together, they further optimize the matrix properties. A high-temperature tensile test was conducted at 800℃ according to GB / T 228.2-2015 standard, and the high-temperature tensile strength reached 880 MPa. This result demonstrates that by adjusting the matrix composition, the final high-temperature mechanical properties of the composite material can be effectively controlled even without changing the type, content, or preparation process of the reinforcing phase.

[0140] Comparative Example 1

[0141] A method for preparing a ceramic-reinforced high-entropy alloy material, which differs from Example 1 only in the ceramic-reinforced phase, that is, in step (3), the composite ceramic particles obtained in Example 1 are replaced with Al2O3 powder (irregular particles with a particle size of 10-45μm, added at 5wt.%).

[0142] Performance results:

[0143] The density of the ceramic-reinforced high-entropy alloy material prepared in this comparative example reached 96.8%. Due to the poor flowability of Al2O3 powder (>45s / 50g), its wettability with the high-entropy alloy matrix was extremely poor, and severe segregation occurred after premixing with the high-entropy alloy matrix powder. The interface between Al2O3 and the high-entropy alloy matrix was clear, without a nanoscale reaction interface layer, and exhibited mechanical bonding. Under thermal stress, severe interface debonding occurred. The high-temperature tensile strength test conducted at 800℃ according to GB / T 228.2-2015 standard showed a high strength of 580MPa, even lower than that of the unreinforced matrix (620MPa).

[0144] Comparative Example 2

[0145] A method for preparing a ceramic-reinforced high-entropy alloy material, which differs from Example 1 only in the ceramic-reinforced phase, that is, in step (3), the composite ceramic particles obtained in Example 1 are replaced with WC powder (irregular particles with a particle size of 10-45μm, added at 5wt.%).

[0146] Performance results:

[0147] The density of the ceramic-reinforced high-entropy alloy material prepared in this comparative example reached 98.2%. Although WC powder has high hardness, its coefficient of thermal expansion with the matrix is ​​large. WC partially melts under high-temperature laser conditions and reacts with elements such as Cr and Co in the matrix to form brittle complex carbides such as (W,Cr,Co)₂C, forming a continuous brittle interface. This interface becomes a rapid crack propagation channel under stress. The high-temperature tensile strength test conducted at 800℃ according to GB / T 228.2-2015 standard showed a high tensile strength of 660 MPa, with limited improvement and exhibiting brittle fracture.

[0148] Comparative Example 3

[0149] A method for preparing a ceramic-reinforced high-entropy alloy material, which differs from Example 2 only in the ceramic-reinforced phase, that is, in step (3), the composite ceramic particles obtained in Example 1 are replaced with CrC powder (irregular particles with a particle size of 10-45μm, added at 15wt.%).

[0150] Performance results:

[0151] The ceramic-reinforced high-entropy alloy material prepared in this comparative example exhibited a small number of pores in its deposited layer. Furthermore, under the same high-temperature creep rupture test conditions as Example 2, its lifespan was only 42 hours. The fracture occurred due to interfacial degradation at the reinforced phase aggregation site. This is because CrC suffers from insufficient thermodynamic stability under high-temperature laser irradiation and subsequent service, leading to partial decomposition and carbon diffusion, resulting in fluctuations in the interfacial layer composition and uneven thickness. In contrast, Example 2, through the high-temperature stability of the composite ceramic phase and the in-situ formation of the nanocomposite interfacial layer, achieved long-term stability in microstructure and properties, resulting in an order-of-magnitude improvement in high-temperature creep rupture lifespan.

[0152] Comparative Example 4

[0153] A method for preparing a ceramic-reinforced high-entropy alloy material differs from Example 2 in that the ceramic-reinforced phase is different. In step (3), the composite ceramic particles obtained in Example 2 are replaced with Al2O3 powder (irregular particles with a particle size of 10-45μm), and the amount of Al2O3 powder added is increased to 25wt.

[0154] During the preparation process, the mixed powder, after premixing Al2O3 powder and high-entropy matrix alloy powder, has almost no flowability, making stable powder feeding and cladding impossible. This inability to achieve stable powder feeding and forming results in difficulty in forming shaped parts. In contrast, the embodiments of this application utilize the synergistic complementarity of multiple ceramic materials and impart excellent flowability to the composite ceramic particles through spheroidization treatment, solving the processability dilemma caused by the poor physical properties of traditional reinforcing phases. Furthermore, through an optimized microstructure nanocomposite interface layer, it overcomes the manufacturing bottleneck of traditional high-ceramic-content metal matrix composites being "unmanufacturable or poorly manufactured."

[0155] Comparative Example 5

[0156] A method for preparing a ceramic-reinforced high-entropy alloy material differs from Example 1 only in that the ceramic reinforcing phase is different. The composite ceramic particles are prepared using the same method as in Example 1, but the composite ceramic particles of this comparative example are composed of the following components by mass percentage: BN: 42.3%, ZrB2: 28.8%, NbN: 15.4%, TiB2: 13.5%.

[0157] Performance results:

[0158] Mechanical property tests showed that the performance of the material obtained in this comparative example was significantly inferior to that of Example 1. Its high-temperature tensile strength, tested at 800°C according to GB / T 228.2-2015 standard, decreased to 650 MPa or lower, a decrease of more than 23% compared to Example 1 (850 MPa).

[0159] The microscopic mechanism lies in the removal of the core load-bearing skeleton in the composite material due to the absence of TiC and TiN. Although phases such as BN and ZrB2 have certain strengthening effects, their hardness and modulus are much lower than those of TiC / TiN, resulting in a significant reduction in their ability to hinder dislocation movement and crack propagation (pinning effect). Furthermore, BN and ZrB2 have different interfacial reactivity with the high-entropy alloy melt, making it difficult to form a strong and tough (Ti,W,Nb)C gradient interfacial layer as in Example 1. This leads to a decrease in interfacial bonding strength, making them prone to becoming microcrack initiation points, all of which contribute to the premature failure of the material under high-temperature loads.

[0160] Comparative Example 6

[0161] A method for preparing a ceramic-reinforced high-entropy alloy material differs from Example 1 only in that the ceramic reinforcing phase is different. The composite ceramic particles are prepared using the same method as in Example 1, but the composite ceramic particles of this comparative example are composed of the following components by mass percentage: TiC: 31%, TiN: 35%, TiB2: 13%, BN: 21%.

[0162] Performance results:

[0163] Mechanical property tests showed that the mechanical properties of the material obtained in this comparative example were between those of Example 1 and Comparative Example 5, but were still significantly lower than those of Example 1. Its high-temperature tensile strength, tested at 800°C according to GB / T 228.2-2015 standard, was approximately 700 MPa, a decrease of about 17.6% compared to Example 1 (850 MPa).

[0164] The microscopic mechanism is as follows: Although TiC (31%) and TiN (35%) are retained in the composite ceramic particles as core reinforcing phases, their combined content (66%) can form a basic load-bearing skeleton and provide a certain dislocation pinning effect. However, the high content of BN phase (16%) and TiB2 phase (18%) has a negative impact on the material properties. The relatively high content of BN is still prone to agglomeration at the interface, and its interfacial reaction products with the high-entropy alloy melt are not as strong and tough as the (Ti,W,Nb)C gradient layer. At the same time, the increased TiB2 may also compete with other components during the reaction process, interfering with the full formation of an ideal, continuous and tough interfacial reaction layer. Ultimately, the interfacial bonding strength between the ceramic phase and the high-entropy alloy matrix is ​​weakened. Under high-temperature loads, these relatively weak interfacial regions become preferred sites for the initiation and propagation of microcracks, thereby reducing the overall load-bearing capacity of the material and causing a significant decrease in its high-temperature strength.

[0165] This result further demonstrates that in the design of ceramic reinforcing phases, it is not only necessary to ensure the sufficient presence of core hard phases such as TiC / TiN, but also to strictly control the content of other phases such as BN that have adverse effects on interfacial bonding in order to obtain the optimal comprehensive performance.

[0166] Although this application has been described in detail with reference to the foregoing embodiments, those skilled in the art should understand that modifications can still be made to the technical solutions described in the foregoing embodiments, or equivalent substitutions can be made to some of the technical features; and these modifications or substitutions do not cause the essence of the corresponding technical solutions to deviate from the spirit and scope of the technical solutions of the embodiments of this application.

Claims

1. A composite ceramic particle-reinforced high-entropy alloy matrix composite material, characterized in that, It includes a high-entropy alloy matrix and composite ceramic particles dispersed in the high-entropy alloy matrix; the particle size of the composite ceramic particles ranges from nanometer to micrometer. The composite ceramic particles and the high-entropy alloy matrix have a nanocomposite interface layer, which is generated in situ by bidirectional interdiffusion of elements between the high-entropy alloy matrix and the composite ceramic particles. The composite ceramic particles comprise the following components by mass percentage: TiC: 28%-30%, TiN: 30%-35%, BN: 10%-13%, ZrB2: 7%-8%, NbN: 3%-5%, TiB2: 9%-17%, and the high-entropy alloy matrix comprises any one of the following multi-element alloys: NiCoCrWMoFe, NiCoCrWMoAl, and NiCoCrWFeMnSi.

2. The composite ceramic particle-reinforced high-entropy alloy matrix composite material according to claim 1, characterized in that, The composite ceramic particles account for 5-30% of the mass fraction of the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

3. The composite ceramic particle-reinforced high-entropy alloy matrix composite material according to claim 1, characterized in that, The loose packing density of the composite ceramic particles is 0.6-4.2 g / cm³. 3 The flow rate is 10-40 s / 50 g.

4. The composite ceramic particle-reinforced high-entropy alloy matrix composite material according to claim 1, characterized in that, The thickness of the nanocomposite interface layer is 40-300 nm.

5. A method for preparing the composite ceramic particle-reinforced high-entropy alloy matrix composite material as described in claim 1, characterized in that, Includes the following steps: Step S1: Mix various ceramic materials such as TiC, TiN, ZrB2, TiB2, BN, and NbN with water and binder to form a slurry; Step S2: The slurry is subjected to spray granulation, high-temperature sintering, and plasma spheroidization in sequence, and then screened to obtain the composite ceramic particles. Step S3: The composite ceramic particles are mixed with high-entropy alloy matrix powder and then subjected to laser additive manufacturing to obtain the composite ceramic particle-reinforced high-entropy alloy matrix composite material.

6. The preparation method according to claim 5, characterized in that, The laser additive manufacturing process employs selective laser melting or laser melting deposition. The parameters of the laser melting deposition process are set as follows: laser power 400-1500W, scanning speed 400-1200mm / min, powder feeding tray rotation speed 800-2800r / min, carrier gas flow rate 3-10L / min, and overlap rate 40-50%. The parameters of the laser selective melting process are set as follows: laser power 200-350W, scanning speed 800-1500mm / s, scanning spacing 0.1-0.13mm, auxiliary powder layer thickness 30-50μm, and overlap rate 40-50%.

7. The preparation method according to claim 5, characterized in that, In step S3, after the laser additive manufacturing, the process further includes: subjecting the high-entropy alloy-based composite material after laser additive manufacturing to solution treatment and aging treatment under vacuum conditions in sequence; In the solution treatment, the high-entropy alloy-based composite material is heated to 1250±10℃, held at that temperature for 2±0.25h, and then cooled to room temperature under an inert atmosphere. In the aging treatment, the high-entropy alloy-based composite material after solution treatment is heated to 750±20℃, held at that temperature for 1±0.25h, and then cooled to room temperature by filling with an inert atmosphere.

8. The preparation method according to claim 5, characterized in that, In step S2, the spray granulation includes: atomizing the slurry under a temperature field of 250-300°C inlet temperature and 100-105°C outlet temperature, and then drying the atomized droplets to form agglomerates; The high-temperature sintering process includes: holding the spray-granulated agglomerates at 750-850℃ for 5-8 hours under a protective atmosphere, and then holding them at 1000-1400℃ for 6-8 hours to obtain the sintered product. The parameters for the plasma spheroidization process include: controlling the plasma beam current to 620-630A, the main gas flow rate to 120-130 SCFH, the auxiliary gas flow rate to 10-11 SCFH, the carrier gas flow rate to 9-12 SCFH, and the powder feeding rate to 4.0-6.0 r / min.

9. The preparation method according to claim 5, characterized in that, The high-entropy alloy matrix powder is prepared by a method comprising the following steps: According to the target composition of the high-entropy alloy matrix, each elemental metal raw material is weighed, and the weighed elemental metal raw materials are subjected to vacuum induction melting under a protective atmosphere to obtain the alloy melt. The alloy melt is atomized under vacuum conditions to obtain spherical or near-spherical high-entropy alloy powder using high-pressure inert gas as the atomization medium. The high-entropy alloy powder is sieved to obtain high-entropy alloy powder of the target particle size as high-entropy alloy matrix powder.

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