High-entropy cerate ceramic and preparation method thereof

Abstract

The invention discloses a high-entropy cerate ceramic and a preparation method thereof, so that on one hand, the chemical composition of the high-entropy cerate ceramic provided by the invention is represented by the following general formula: (A < 0.2 > B < 0.2 > Nd < 0.2 > Gd < 0.2 > Tm < 0.2 >) 2Ce2O7, and the chemical composition of the high-entropy cerate ceramic is represented by the following general formula: (A < 0.2 > B < 0.2 > Nd < 0.2 > Gd < 0.2 > Tm < 0.2 >) 2Ce2O7, wherein A and B are respectively and independently selected from any one of Y, La, Pr and Sm, and A and B are different. The prepared high-entropy cerate ceramic material has low thermal conductivity, high thermal expansion coefficient and excellent high-temperature phase stability, and can meet the performance requirements of thermal barrier coatings of next-generation aero-engines.

Description

Technical Field

[0001] This invention belongs to the field of high-temperature protective materials technology, and particularly relates to a high-entropy cerate ceramic and its preparation method. Background Technology

[0002] Thermal barrier coatings (TBCs) are a key technology for improving the service temperature and lifespan of hot-end components such as aero-engines and gas turbines. Currently, 6–8 wt.% Y₂O₃-stabilized zirconium oxide (YSZ) is widely used, exhibiting low thermal conductivity, high coefficient of thermal expansion, and good fracture toughness below 1200℃. However, as engine inlet temperatures increase to above 1500℃, YSZ undergoes a harmful phase transition from tetragonal to monoclinic at high temperatures. This phase change, accompanied by volume changes, leads to coating cracking and peeling, failing to meet the requirements of next-generation power systems.

[0003] Rare earth cerates are considered potential high-temperature thermal barrier coating materials due to their defective fluorite structure, high oxygen vacancy concentration, low thermal conductivity, and high melting point. However, pure cerates are prone to reduction at high temperatures, leading to lattice instability and decreased sintering resistance. High-entropy ceramics, through multi-principal element design, can form a single solid solution phase, exhibiting significant lattice distortion and slow diffusion effects, which helps to further improve the thermophysical properties and high-temperature stability of the material. However, compared with high-entropy zirconate and hafnium salts, high-entropy cerate ceramics still have problems such as higher thermal conductivity, lower coefficient of thermal expansion, and insufficient ultra-high temperature stability. Summary of the Invention

[0004] The purpose of this invention is to provide a high-entropy cerate ceramic material that combines low thermal conductivity, high coefficient of thermal expansion and excellent high-temperature phase stability, and to provide a reproducible synthesis method suitable for large-scale preparation, so as to meet the performance requirements of thermal barrier coatings for next-generation aero-engines.

[0005] To this end, the present invention provides a high-entropy cerate ceramic, the chemical composition of which is represented by the following general formula: (A 0.2 B 0.2 Nd 0.2 Gd 0.2 Tm 0.2 )2Ce2O7; where, A and B are each independently selected from any one of Y, La, Pr, and Sm, and A and B are not the same.

[0006] Specifically, the ceramic has a single-phase defect fluorite structure and a thermal conductivity of less than 1.70 W·m at 1000℃. -1 ·K -1 The coefficient of thermal expansion is higher than 12.5 × 10¹⁰. -6 / K remains structurally stable after being heat-treated at 1600℃ for 100 hours.

[0007] Specifically, A and B are selected from any of the following element pairs: Pr and Y, or La and Sm, or Sm and Y.

[0008] In another aspect, this invention provides a method for preparing high-entropy cerate ceramics, comprising the following steps: S1. Raw material pretreatment: All raw material powders, including Ce2O3, Tm2O3, Nd2O3, Gd2O3 and the corresponding oxides of elements A and B, are pre-calcined. S2. Ingredients and mixing: Weigh each raw material powder after step S1 according to the stoichiometric ratio of the general formula described in claim 1, and mix it together with the ball milling media and dispersant by ball milling to obtain a uniform slurry; S3. Drying and pre-firing: The slurry is dried to obtain a mixed powder, and then sintered for the first time at a temperature of 1300℃ to 1400℃ to obtain high-entropy ceramic initial powder; S4. Secondary processing and molding: The high-entropy ceramic powder is subjected to secondary ball milling, drying, and sieving, then a binder is added, and the ceramic green body is obtained by dry pressing. S5. Final sintering: After debinding the ceramic green body, a second sintering is carried out in an air atmosphere at 1450°C to 1550°C to obtain the high-entropy cerate ceramic bulk material.

[0009] Specifically, in step S2, the ball milling speed is 200 rpm to 300 rpm, and the ball milling time is 2 hours to 4 hours.

[0010] Before final sintering, the powder undergoes a first sintering process to address structural and performance defects in the ball-milled powder, achieving component homogenization, phase structure stabilization, and densification of the high-entropy ceramic, thus obtaining high-entropy ceramic raw powder that meets application requirements. The first sintering temperature is controlled at 1300℃ to 1400℃ (lower than the second sintering temperature (1450℃ to 1550℃)). This design aims to achieve the formation of the main phase, particle neck bonding, and powder densification under relatively mild conditions. The second sintering temperature is increased to 1450℃ to 1550℃, utilizing the higher temperature to enhance oxygen ion and cation diffusion, promote the further dissolution of residual impurities into the crystal lattice, eliminate closed pores, and improve density. Simultaneously, the grain boundary structure and stress state are optimized, ultimately yielding high-entropy oxide ceramics with uniform phase, high density, and controllable microstructure, providing structural assurance for achieving excellent mechanical, thermal, and high-temperature stability properties.

[0011] Specifically, in step S3, the holding time for the first sintering is 1 to 3 hours.

[0012] Specifically, in step S5, the holding time for the second sintering is 4 to 8 hours.

[0013] Specifically, in step S5, the heating procedure for the debinding process and the second sintering is as follows: first, the temperature is raised to 800°C at a first heating rate and held at that temperature, and then the temperature is raised to the target sintering temperature at a second heating rate lower than the first heating rate.

[0014] The above-mentioned two-step heating regime of 800℃ debinding and 1500℃ sintering: The first step is to perform isothermal holding at 800℃, which can not only slowly and thoroughly remove the binder and other organic components introduced during the ball milling process, avoiding cracking, pore defects and residual carbon contamination of the green body caused by rapid gas release, but more importantly, it can achieve low-temperature pre-solid phase reaction and preliminary solid solution of multi-component oxides at this temperature, forming a precursor phase with uniform composition, effectively eliminating the local composition fluctuations and chemical potential gradients inside the powder after ball milling, and providing an initial green body with uniform structure and composition for subsequent high-temperature sintering. The second step employs a slower heating rate to 1500℃. On one hand, controlling the heating rate suppresses the low-melting-point eutectic liquid phase induced by the rapid migration of local components and limits the grain boundary migration rate, preventing abnormal grain growth and ensuring the uniformity of the microstructure. On the other hand, the slow heating process continuously enhances the grain boundary and lattice diffusion of cations and oxygen ions, promoting the growth of intergranular necks, spheroidization of pores, and shrinkage and discharge, achieving gradual densification. At the same time, it promotes the complete dissolution of residual impurities into the lattice and ultimately forms a single-phase high-entropy solid solution, thereby synergistically obtaining a microstructure with high density, pure phase, and uniform grain size.

[0015] Specifically, the first heating rate is controlled at 4-5 ℃ / min; the second heating rate is controlled at 2-3 ℃ / min.

[0016] This application employs a lattice distortion engineering strategy to maximize the difference in ionic radii, with Nd 3+ Gd 3+ Tm 3+ Using Nd as the core, an ion radius spectrum covering large, medium, and small scales was constructed, and La, Sm, Pr, and Y rare earth ions were synergistically introduced to form a high-entropy solid solution design with multi-size distortion, multi-electronotropic regulation, and multi-type defect coupling. Among them, large-size Nd 3+ With small size Tm 3+ The formation of extreme size contrasts can generate significant lattice strain fields, leading to a reduction in thermal conductivity through strong phonon scattering. This is particularly relevant for Gd crystals with intermediate sizes. 3+ This can buffer extreme distortion stress, preventing the distortion from exceeding the ±15% tolerance of the fluorite structure and causing phase separation. Furthermore, the size disorder parameter of the pentagonal combination is significantly higher than that of binary or ternary systems, ensuring that the lattice thermal conductivity is reduced to the theoretical minimum. 3+ With Nd 3+ Forming a super-large ion terminus, with Y 3+ Tm 3+Small ion ends create extreme size contrasts, enhancing lattice distortion and phonon scattering; Sm 3+ Y 3+ Fill in the intermediate transition interval to achieve a continuous and smooth radius distribution; Gd 3+ As core buffer ions, they work in conjunction with Sm and Y to control the lattice strain within the tolerance of the fluorite structure, thus avoiding phase separation. Regarding the electronegativity gradient, the low electronegativity of La, Nd, and Pr synergistically suppresses Ce. 4+ Reduction, high electronegativity of Tm and Y enhances lattice oxygen stability, Gd and Sm balance charge distribution, and Pr... 3+ / Pr 4+ The valence state is dynamically adjusted to achieve dual stability of valence state and structure. In defect engineering, the total vacancy concentration is precisely controlled below the critical threshold by combining low, medium, and high oxygen vacancy forming ions. This utilizes defect scattering to reduce thermal conductivity while suppressing high-temperature vacancy migration and ordering. Regarding thermal expansion, the CTE is adjusted to be nearly matched with the high-temperature alloy substrate through large / small ion reverse matching, and the high-entropy anharmonic effect is combined to alleviate thermal stress mismatch. This multi-rare-earth synergistic strategy maximizes lattice distortion and minimizes thermal conductivity while strictly ensuring single-phase solid solution stability at 1500–1600℃, achieving a synergistic improvement in low thermal conductivity, high structural stability, and excellent thermal matching.

[0017] By designing electronegativity gradients of rare earth elements, synergistic regulation of electronegativity and valence stability is achieved, resulting in complementary optimization of electronic structure: low electronegativity Nd 3+ It can enhance the ionicity of Ce-O bonds and effectively suppress Ce under high-temperature reducing atmospheres. 4+ To Ce 3 + The transformation avoids structural collapse caused by excessive oxygen vacancy generation; the high electronegativity of Tm 3+ It can strengthen the binding energy of lattice oxygen and improve the chemical stability of the fluorite phase, especially ensuring the integrity of the crystal structure under low oxygen partial pressure; while Gd, with its intermediate electronegativity, can enhance the binding energy of lattice oxygen and improve the chemical stability of the fluorite phase. 3+ It can balance the overall electron cloud distribution of the system, avoid the problem of cation ordering caused by local charge mismatch, and maintain the structural and valence stability of the solid solution in all aspects.

[0018] Using trivalent rare earth elements to replace Ce 4+ The generated charge compensates for oxygen vacancies, and defect engineering optimization is achieved by leveraging the differences in the formation energy of oxygen vacancies in different rare earth elements: large-radius Nd 3+ Oxygen vacancies have low formation energy and are easily generated, but an excess of them can lead to abnormal thermal expansion and structural relaxation. Therefore, Gd with a moderate formation energy is used to address this issue. 3+ Tm has high formation energy and can stabilize its lattice with a small radius. 3+Synergistic regulation precisely controls the total oxygen vacancy concentration below the critical threshold, which reduces thermal conductivity by relying on the scattering effect of oxygen vacancies and effectively avoids the orderly aggregation of vacancies. It can significantly suppress vacancy migration and lattice oxygen loss in high-temperature environments above 1500℃, ensuring the long-term structural stability of the material during service.

[0019] Inverse matching optimization of thermal expansion coefficient through ionic radius weighting effect: This involves optimizing the thermal expansion coefficient of large-radius Nd with high CTE contribution. 3+ With low CTE contribution and small radius Tm 3+ Equimolar mixing modulates the theoretical CTE of the system to 11-13×10⁻⁶. -6 / K, the CTE of this pentagonal combination is closer to that of the alloy substrate; at the same time, the high entropy effect of the system enhances the lattice anharmonicity, which can improve the CTE to a certain extent at high temperatures and effectively alleviate the thermal stress mismatch problem between the coating and the substrate.

[0020] The five rare earth elements selected in this system exhibit differentiated ionic radius distributions. The size disorder effect resulting from moderate radius differences is key to reducing thermal conductivity. This achieves low thermal conductivity by enhancing phonon scattering through lattice distortion, while controlling the radius differences within the Hume-Rothery solid solution limit to avoid phase separation. Simultaneously, this rare earth combination has a low enthalpy of formation, resulting in more stable chemical bonding between rare earth ions and the CeO2 matrix. This leads to a more negative formation free energy ΔG in the solid solution, effectively suppressing cation segregation and phase separation during high-temperature sintering at 1500-1600℃ and in service, ensuring the long-term stability of the fluorite structure.

[0021] Compared with the prior art, the present invention has the following beneficial effects: 1. Through a high-entropy design using five rare earth elements (including Tm, Nd, Gd, A, B) and Ce, while maintaining low thermal conductivity (<1.70 W·m), -1 ·K -1 While achieving a high coefficient of thermal expansion (>12.5×10 at 1000℃), it also obtained a high coefficient of thermal expansion (>12.5×10). -6 / K), which has better compatibility with high-temperature alloy substrates.

[0022] 2. High-entropy cerates maintain a single-phase defect fluorite structure at high temperatures of 1500–1600℃ for a long time without harmful phase transformations or decomposition, and their structural stability is significantly better than that of traditional rare earth cerates.

[0023] 3. It adopts the classic solid-state reaction method and ball milling process, the raw materials are readily available, the process parameters are wide-ranging, and it is easy to achieve precise control of composition and structure, making it suitable for large-scale production. Attached Figure Description

[0024] To more clearly illustrate the technical solutions in the embodiments of the present invention, the accompanying drawings used in the description of the embodiments will be briefly introduced below. Obviously, the accompanying drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0025] Figure 1 The images show the XRD, particle size distribution, SEM, and corresponding EDS images of the high-entropy ceramic powder prepared in Example 1. Figure 2 These are EDS, high-resolution transmission electron microscopy, and selected area electron diffraction patterns of individual ceramic particles of the high-entropy ceramic powder particles prepared in Example 1. Figure 3 This is a SEM image of the high-entropy ceramic prepared in Example 1, along with its grain size distribution. Figure 4. XRD patterns of the high-entropy ceramic prepared in Example 1 before and after the microstructure stability test at 1600℃ for 100 hours; Figure 5 Thermal conductivity diagram of the high-entropy ceramic prepared in Example 1. Detailed Implementation

[0026] The technical solutions of the embodiments of the present invention will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0027] The present invention provides a high-entropy cerate ceramic, the chemical composition of which is represented by the following general formula: (A) 0.2 B 0.2 Nd 0.2 Gd 0.2 Tm 0.2 )2Ce2O7; where, A and B are each independently selected from any one of Y, La, Pr, and Sm, and A and B are not the same. The ceramic has a single-phase defect fluorite structure and a thermal conductivity of less than 1.70 W·m at 1000℃. -1 ·K -1 The coefficient of thermal expansion is higher than 12.5 × 10⁻⁶. -6 / K remains structurally stable after being heat-treated at 1600℃ for 100 hours.

[0028] The raw materials for high-entropy cerate ceramics are Ce₂O₃, Tm₂O₃, Nd₂O₃, Gd₂O₃ and La₂O₃, Sm₂O₃, Y₂O₃, Pr₆O₃. 11Two rare earth oxides (purity ≥ 99%) were used in the process, with the oxides being oxidized at 1300°C to 1400°C and 1450°C to 1550°C. Example 1 A high-entropy cerate ceramic is prepared by the following method: Take Pr6O respectively 11 Y2O3, Tm2O3, Nd2O3, Gd2O3, and Ce2O3 powders were calcined at 1000℃ for 2 hours to remove impurities such as carbon dioxide and moisture adsorbed in the powder. Pr6O was weighed according to a molar ratio of 0.4:0.4:0.4:0.4:0.4:2. 11 Y2O3, Tm2O3, Nd2O3, Gd2O3, Ce2O3 powders were used to obtain powder raw materials; Anhydrous ethanol and milling beads were added to the above-mentioned powder raw materials. The mass ratio of powder raw materials, anhydrous ethanol, and milling beads was 3:4:2. Zirconia milling beads were used, the milling speed was 300 rpm / min, and the milling time was 4 hours. The milled slurry was placed in a drying oven and dried at 100 °C for 10 hours. The dried powder was then sintered at 1400 °C for 2 hours. Afterward, the same parameters were used for a second milling. After the milling was completed, the powder was dried and passed through a 100-mesh sieve to obtain the powder. The XRD pattern of the prepared high-entropy cerate ceramic powder is shown below. Figure 1 As shown in (a), the particle size distribution diagram is as follows: Figure 1 As shown in (b), the SEM image and the corresponding EDS image are as follows: Figure 1 As shown in (c), the EDS image of the particles is as follows: Figure 2 As shown in (a), the high-resolution transmission image of the powder is as follows: Figure 2 As shown in (b); the selected area electron diffraction pattern of a single ceramic particle is as follows: Figure 2 As shown in (c), from the attached Figure 1 As can be seen from XRD, the product shows a single defect fluorite structure, while SEM and EDS indicate uniform elemental distribution (see attached diagram). Figure 1 c, 2a), TEM and SAED confirmed its crystalline characteristics (see appendix). Figure 2 b, c).

[0029] PVA was added to the obtained high-entropy cerate ceramic powder and stirred until homogeneous, wherein the mass ratio of PVA powder to deionized water in the PVA solution was 5:95. The mixture was then placed in a pressing mold, subjected to a pressure of 6 t, and held for 5 min to obtain a green body. The green body was placed in a muffle furnace and heated to 800 °C at a rate of 5 °C / min, held for 2 h, then heated to 1480 °C at a rate of 2 °C / min, held for 6 h, and finally cooled in the furnace to obtain high-entropy cerate ceramics. SEM images of the prepared high-entropy cerate ceramic particles are shown below. Figure 3As shown in (a), the grain size distribution diagram is as follows: Figure 3 As shown in (b), from Figure 3 It can be seen that the ceramic block is dense and the grain size is uniformly distributed.

[0030] The ceramic block was placed in a muffle furnace and subjected to a tissue stability test at 1600 °C for 100 h. The XRD patterns of the high-entropy ceramic before and after the 100-hour tissue stability test at 1600 °C are shown below. Figure 4 As shown, from Figure 4 It can be seen that there is no phase change after annealing at 1600℃ for 100 hours, proving its excellent high-temperature phase stability. This is because the rare earth combination has a low formation enthalpy (-0.0166eV / atom), which makes the chemical bonding between rare earth ions and CeO2 matrix more stable, and makes the formation free energy ΔG of solid solution more negative. This effectively suppresses cation segregation and phase separation during high-temperature sintering at 1500-1600℃ and service, ensuring the long-term stability of the fluorite structure.

[0031] The thermal conductivity and coefficient of thermal expansion of the high-entropy ceramic were tested. The thermal conductivity of the high-entropy ceramic was as follows: Figure 5 As shown in the figure, the thermal conductivity test results indicate that the material of this invention has a much lower thermal conductivity than YSZ at high temperatures, demonstrating outstanding thermal insulation performance. It also possesses a high coefficient of thermal expansion (12.92). -6 / K), which is closely matched with the high-temperature alloy substrate. Example

[0032] Unlike Example 1, La2O3, Sm2O3, Tm2O3, Nd2O3, Gd2O3, and Ce2O3 powders were calcined at 1000℃ for 2 hours to remove adsorbed impurities such as carbon dioxide and moisture from the powders. The La2O3, Sm2O3, Tm2O3, Nd2O3, Gd2O3, and Ce2O3 powders were weighed according to a molar ratio of 0.4:0.4:0.4:0.4:0.4:2 to obtain the powder raw material.

[0033] The green body was placed in a muffle furnace and heated to 800 °C at a rate of 4 °C / min, held for 2 h, then heated to 1520 °C at a rate of 2 °C / min, held for 6 h, and finally cooled in the furnace to obtain high-entropy cerate ceramic. Its thermal conductivity, coefficient of thermal expansion, and microstructure stability at 1600 °C were then tested. The microstructure stability test showed that the high-entropy cerate ceramic of this embodiment showed no phase change after annealing at 1600 °C for 100 hours. This is because the rare earth combination also has a low enthalpy of formation (-0.0163 eV / atom). The thermal conductivity test results showed that the material of this invention has a much lower thermal conductivity than YSZ at high temperatures, exhibiting excellent thermal insulation performance. It also has a high coefficient of thermal expansion (12.96). -6 / K), which is closely matched with the high-temperature alloy substrate. Example

[0034] Unlike Example 1, Sm2O3, Y2O3, Tm2O3, Nd2O3, Gd2O3, and Ce2O3 powders were calcined at 1000℃ for 2 hours to remove impurities such as carbon dioxide and moisture adsorbed in the powders; Sm2O3, Y2O3, Tm2O3, Nd2O3, Gd2O3, and Ce2O3 powders were weighed according to a molar ratio of 0.4:0.4:0.4:0.4:0.4:2 to obtain powder raw materials.

[0035] The green body was placed in a muffle furnace and heated to 800 °C at a rate of 5 °C / min, held for 2 h, then heated to 1500 °C at a rate of 2 °C / min, held for 6 h, and finally cooled in the furnace to obtain high-entropy cerate ceramic. Its thermal conductivity and coefficient of thermal expansion were tested. The thermal conductivity and coefficient of thermal expansion are shown in Table 1. The thermal conductivity test results show that the material of this invention has a much lower thermal conductivity than YSZ at high temperatures, exhibiting outstanding thermal insulation performance. The microstructure stability test results show that the high-entropy cerate ceramic of this embodiment showed no phase change after annealing at 1600 °C for 100 hours. This is because this rare earth combination also has a low enthalpy of formation (-0.0149 eV / atom). It also has a high coefficient of thermal expansion (12.88 eV / atom). -6 / K), which is closely matched with the high-temperature alloy substrate.

[0036] Comparative Example 1 Unlike Example 1, Comparative Example 1 was heated directly to 1500 °C at a rate of 5 °C / min and held at that temperature for 6 h. The resulting high-entropy cerate ceramic was then tested for thermal conductivity and coefficient of thermal expansion. The results are shown in Table 1. As can be seen from Table 1, the thermal conductivity of Comparative Example 1 is significantly higher. This is because rapid heating in one step causes the organic decomposition, solid-phase reaction, densification, solid solution, and grain growth processes to be highly coupled and fiercely competitive, resulting in uneven composition, insufficient solid solution, defect clustering, a large number of microcracks and closed pores, and abnormal grain growth. Ultimately, this manifests as decreased high-temperature stability, increased thermal conductivity, and mismatch in the coefficient of thermal expansion. Example 1 uses stepwise slow heating to achieve multi-component high-entropy uniform solid solution, continuous strengthening of the lattice distortion field, controllable oxygen vacancy dispersion, and uniform and refined microstructure through pre-calcination degassing and pre-reaction, segmented diffusion solid solution, progressive densification, and controllable grain growth. This systematically improves the high-temperature structural stability, low thermal conductivity characteristics, and thermal expansion matching.

[0037] Comparative Example 2 Unlike Example 1, Comparative Example 2 used the same first and second sintering temperatures of 1480 °C. The thermal conductivity and coefficient of thermal expansion are shown in Table 1. Table 1 shows that the overall thermophysical and mechanical properties of the material deteriorated significantly. This is because the excessively high pre-sintering temperature after ball milling causes excessive necking and hard agglomeration of the powder particles, leading to a decrease in surface energy and a significant reduction in sintering activity. It also induces premature local solidification of some components and the formation of impurity phases, disrupting the compositional homogeneity of the high-entropy system. Furthermore, high-temperature pre-sintering causes premature coarsening of the powder grains and promotes the excessive generation and clustering of oxygen vacancies induced by low-vacancy-forming-energy rare-earth ions, weakening the phonon scattering effect of defects. These problems directly lead to difficulties in densification, coarse microstructure, insufficient solid solution, and unstable phase structure during the secondary sintering process, ultimately resulting in decreased high-temperature structural stability, increased thermal conductivity, and a mismatch and increased fluctuation in the coefficient of thermal expansion.

[0038] Comparative Example 3 Unlike Example 2, in Comparative Example 3, the temperature was first increased to 600 °C at a rate of 5 °C / min, held for 2 h, and then increased to 1600 °C at a rate of 2 °C / min. Its thermal conductivity and coefficient of thermal expansion are shown in Table 1. As can be seen from Table 1, the thermal conductivity and coefficient of thermal expansion of the material deteriorate accordingly. This is because excessively low debinding temperatures lead to incomplete decomposition of organic components such as binders and dispersants in the green body, resulting in residual carbon and undecomposed organic matter. During subsequent high-temperature sintering, secondary gas generation occurs, inducing closed pores, microcracks, and internal pore defects. Simultaneously, residual carbon triggers a localized carbothermic reduction reaction, leading to Ce... 4+ Excessive reduction leads to premature and excessive generation of oxygen vacancies. If an excessively high final sintering temperature is applied on top of this, it will further exacerbate abnormal grain growth, oxygen vacancy aggregation and ordering, grain boundary embrittlement, and even localized overheating. This disrupts the compositional homogeneity and lattice distortion field continuity of the high-entropy solid solution, leading to the risk of cation segregation and phase separation. The aforementioned defect coupling effects will directly result in a significant decrease in the material's high-temperature structural stability. Simultaneously, weakened grain boundary scattering, reduced defect scattering efficiency, and the formation of localized high thermal conductivity channels cause the thermal conductivity to deviate significantly from the design low value. Furthermore, multiphase coexistence, porosity, cracks, and abnormal lattice distortion collectively cause inaccurate and enhanced nonlinearity in the coefficient of thermal expansion, exacerbating the thermal stress mismatch between the coefficient of thermal expansion and the alloy substrate, ultimately deteriorating the material's overall thermophysical properties.

[0039] Table 1

[0040] As can be seen from Table 1, (A) prepared according to the method of the present invention 0.2 B 0.2 Nd 0.2 Gd 0.2 Tm 0.22Ce₂O₇ high-entropy cerate ceramics have a single-phase defect fluorite-type structure, uniform composition, and low thermal conductivity (1.52 W·m⁻¹) when sintered at 1500 ℃. -6 It exhibits excellent high-temperature thermal stability (K, 1000℃).

[0041] The technical solutions in the embodiments of the present invention have been clearly and completely described. Obviously, the described embodiments are only a part of the embodiments of the present invention, and not all of them. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

Claims

1. A high-entropy cerate ceramic, characterized in that: The chemical composition of the high-entropy cerate ceramic is represented by the following general formula: (A 0.2 B 0.2 Nd 0.2 Gd 0.2 Tm 0.2 )2Ce2O7; where, A and B are each independently selected from any one of Y, La, Pr, and Sm, and A and B are not the same.

2. The high-entropy cerate ceramic according to claim 1, characterized in that: The ceramic has a single-phase defect fluorite structure and a thermal conductivity of less than 1.70 W·m at 1000℃. -1 ·K -1 The coefficient of thermal expansion is higher than 12.5 × 10⁻⁶. -6 / K remains structurally stable after being heat-treated at 1600℃ for 100 hours.

3. The high-entropy cerate ceramic according to claim 1 or 2, characterized in that: A and B are selected from any of the following element pairs: Pr and Y, or La and Sm, or Sm and Y.

4. A method for preparing high-entropy cerate ceramics according to any one of claims 1-3, characterized in that: Includes the following steps: S1. Raw material pretreatment: All raw material powders, including Ce2O3, Tm2O3, Nd2O3, Gd2O3 and the corresponding oxides of elements A and B, are pre-calcined. S2. Ingredients and mixing: Weigh each raw material powder after step S1 according to the stoichiometric ratio of the general formula described in claim 1, and mix it together with the ball milling media and dispersant by ball milling to obtain a uniform slurry; S3. Drying and pre-firing: The slurry is dried to obtain a mixed powder, and then sintered for the first time at a temperature of 1300°C to 1400°C; S4. Secondary processing and molding: The high-entropy ceramic powder is subjected to secondary ball milling, drying, and sieving, then a binder is added, and the ceramic green body is obtained by dry pressing. S5. Final sintering: After debinding the ceramic green body, a second sintering is carried out in an air atmosphere at 1450°C to 1550°C to obtain the high-entropy cerate ceramic bulk material.

5. The preparation method according to claim 4, characterized in that: In step S2, the ball milling speed is 200 rpm to 300 rpm, and the ball milling time is 2 hours to 4 hours.

6. The preparation method according to claim 4, characterized in that: In step S3, the holding time for the first sintering is 1 to 3 hours.

7. The preparation method according to claim 4, characterized in that, In step S5, the holding time for the second sintering is 4 to 8 hours.

8. The preparation method according to claim 4, characterized in that, In step S5, the heating procedure for the debinding process and the second sintering is as follows: first, the temperature is raised to 800°C at a first heating rate and held at that temperature, and then the temperature is raised to the target sintering temperature at a second heating rate lower than the first heating rate.

9. The preparation method according to claim 8, characterized in that: The first heating rate is controlled at 4-5 ℃ / min; the second heating rate is controlled at 2-3 ℃ / min.

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