Ceramic-based PCBN composite material and preparation method thereof

Abstract

The invention provides a ceramic-based PCBN composite material and a preparation method thereof.The preparation method comprises the steps that a ceramic matrix and a titanium source are placed in an inert gas environment to be subjected to ball-milling mixing, cubic boron nitride powder and a functional auxiliary agent are added, a heating reaction is conducted, a composite precursor is obtained, the functional auxiliary agent comprises an amorphous carbon source, and the amorphous carbon source is added into the ceramic matrix to be subjected to ball-milling mixing; the composite precursor is loaded into a forming die, pre-compression is applied for pre-sintering, the amorphous carbon source is cracked, a nano reaction layer is generated between the cubic boron nitride powder and the ceramic matrix, and a pre-sintered body is obtained; and sintering the pre-sintered body under a high-temperature condition, and raising the ambient pressure and temperature to obtain the ceramic-based PCBN composite material. The steps respectively provide a continuous process of interface coating, interface generation and interface densification, so that the final ceramic-based PCBN composite material can effectively overcome the defect of poor interface bonding.

Description

Technical Field

[0001] This invention belongs to the field of ceramic composite materials technology, and particularly relates to a ceramic-based PCBN composite material and its preparation method. Background Technology

[0002] Polycrystalline cubic boron nitride (PCBN) composites are a class of high-performance composite materials formed by sintering cubic boron nitride (cBN) as the main superhard phase with ceramic or metallic binder phases. Their core characteristics lie in the following: The material's interior is composed of numerous fine cBN grains that, under high temperature and high pressure (HPHT) conditions, form a continuous and dense superhard framework structure through grain boundary reconstruction, sintering neck connections, and interfacial reaction layer reinforcement. Simultaneously, a ceramic phase acts as a binder phase, filling the spaces between the cBN grains to construct a composite network structure with high interfacial bonding capacity and good toughness, resulting in extremely high hardness, wear resistance, and thermal stability. PCBN can achieve a hardness of 40–60 GPa, approaching that of diamond, while possessing significantly higher chemical inertness, oxidation resistance, and thermal shock resistance than diamond. It excels particularly in high-wear machining applications such as cutting hardened steel, heat-resistant alloys, sintered iron, and nickel-based high-temperature alloys, making it a key material in machining, precision manufacturing, and the production of equipment for extreme working conditions.

[0003] In traditional ceramic-based PCBN composites, the ceramic matrix and cubic boron nitride (cBN) belong to two material systems with drastically different crystal structures and significantly different coefficients of thermal expansion. This leads to extremely high residual thermal stress at the interface during high-temperature sintering and cooling. Furthermore, the weak chemical affinity between the ceramic matrix and cBN results in a lack of effective chemical bonds at the interface, often leading to a brittle, crack-prone, and stress-poorly bonded state. Under traditional sintering processes, the absence of a dedicated transition phase or reaction layer to alleviate stress mismatch directly causes problems such as product cracking, low transverse fracture strength, and reduced wear resistance. Summary of the Invention

[0004] The technical problem to be solved by the present invention is to provide a ceramic-based PCBN composite material and its preparation method, which aims to solve the problem of poor mechanical properties and insufficient wear resistance of PCBN composite materials caused by poor interfacial bonding.

[0005] To solve the above-mentioned technical problems, the present invention is implemented as follows: The present invention proposes a method for preparing a ceramic-based PCBN composite material, the steps of which include: S1. Place the ceramic matrix and titanium source in an inert gas environment, ball mill and mix them, then add cubic boron nitride powder and functional additives, and heat to react to obtain a composite precursor, wherein the functional additives include an amorphous carbon source. S2. The composite precursor is loaded into the forming mold and pre-pressure is applied for pre-sintering, so that the amorphous carbon source is decomposed and a nano-reaction layer is generated between the cubic boron nitride powder and the ceramic matrix to obtain the pre-sintered body. S3. The pre-sintered body is first sintered under high temperature conditions, and then the ambient air pressure and temperature are increased to obtain ceramic matrix PCBN composite material.

[0006] In some embodiments of the present invention, the ceramic matrix includes SiAlON powder and alumina micro powder, and step S1 includes: S1.1 Add SiAlON powder, alumina micro powder and titanium source powder to a ball mill jar in a mass ratio of (40-60):(10-20):(10-20), introduce argon or nitrogen to form an inert protective atmosphere, and use a planetary ball mill at 200-250 rpm for 4-6 hours to obtain mixed powder. S1.2 Add cubic boron nitride powder to a ball mill jar at 15-30 wt% of the mixed powder mass, maintain an inert protective atmosphere of argon or nitrogen, and continue ball milling for 1-2 hours at 150-180 rpm to obtain a composite mixed powder containing an ultrahard phase. S1.3 Add the functional additives to the ball mill jar at 3-10 wt% of the composite powder mass, and heat and stir at 80-120℃ to obtain a composite powder with amorphous carbon precursor coated on the surface. S1.4 Place the composite powder coated with amorphous carbon precursor in a tube furnace under an inert atmosphere and heat it at 200-300℃ for 1-2 hours to obtain the composite precursor.

[0007] In some embodiments of the present invention, the functional additives include an amorphous carbon source and an interface activator. The amorphous carbon source includes at least one of phenolic resin, polyimide, carbon black powder, glucose or sucrose solid carbon source. The interface activator includes at least one of lithium nitride, magnesium nitride, lithium borate, calcium fluoride, lithium hydroxide, and lithium nitrate. The titanium source includes at least one of titanium nitride, titanium carbide, and titanium carbonitride.

[0008] In some embodiments of the present invention, step S1 further includes the following step before step S1.2: Cubic boron nitride powder is evenly spread in a ceramic boat, placed in a tube furnace with an inert atmosphere, and heated at 500-600°C for 0.5-1 hour with a continuous flow of pure nitrogen or argon.

[0009] In some embodiments of the present invention, step S1 further includes the following step before step S1.2: Cubic boron nitride powder is placed in a low-pressure plasma reaction chamber, and nitrogen or argon gas is introduced at a radio frequency power of 50-150W to form a plasma state. The processing time is 5-15 minutes.

[0010] In some embodiments of the present invention, step S1 further includes the following step before step S1.2: The cubic boron nitride powder and ceramic grinding beads were placed together in a ball mill and ball-milled at a speed of 300-400 rpm for 10-30 minutes.

[0011] In some embodiments of the present invention, step S2 includes: S2.1. Fill the composite precursor into the silicon nitride mold and cold press it for 1 to 3 minutes with an initial forming pressure of 5 to 10 MPa to obtain the pressed blank. S2.2 Place the pressed blank into a tube furnace and heat it at 700-900℃ for 10-60 minutes in an inert atmosphere of nitrogen, argon or helium to cause thermal decomposition of the amorphous carbon precursor and in-situ deposition of an amorphous carbon nano-reaction layer at the interface of SiAlON, alumina and cubic boron nitride grains. S2.3. The pressed blank with the amorphous carbon nano-reaction layer is kept at 700-900℃ and then subjected to a pressure of 5-10MPa for 5-15 minutes. After releasing the pressure, it is kept at 700-900℃ for 10-30 minutes to obtain the pre-sintered body.

[0012] In some embodiments of the present invention, step S3 includes: S3.1. Load the pre-sintered body into the vacuum sintering furnace and evacuate to 10°C. -2 ~10 -3 Pa, then the temperature is raised to 1300-1500℃ and held for 30-90 minutes to obtain the preform; S3.2. The precast body is placed into the ultra-high pressure cubic pressure chamber and heated to 1500-1700℃ under a pressure of 4-6 GPa, and kept at that temperature for 5-20 minutes. S3.3 Gradually release the pressure of 4-6 GPa and slowly cool it to room temperature at a rate of 5-20℃ / min to obtain ceramic matrix PCBN composite material.

[0013] This invention proposes a ceramic-based PCBN composite material, which is prepared by the method described above.

[0014] Compared with existing technologies, the ceramic-based PCBN composite material and its preparation method disclosed in this invention have the following advantages: In step S1, the ceramic matrix and titanium source are ball-milled and mixed in an inert gas environment, causing the ceramic matrix to generate a high-energy activated surface under mechanical force. Then, cubic boron nitride powder is uniformly introduced and pre-coated onto the particle surface by an amorphous carbon source in the functional additive under heating conditions, thus establishing a reactive and adhesive interfacial precursor layer between the ceramic phase and the cBN phase. The technical feature of step S2 is that pre-pressure is applied to the composite precursor and thermally decomposed, causing the amorphous carbon source to generate a continuous nano-reactive layer in situ between the cBN and the ceramic matrix. This carbon layer can exist as an interfacial transition phase during subsequent sintering, reducing direct contact and mitigating stress concentration between the two materials. The technical feature of step S3 is that the ceramic matrix is ​​first initially densified at high temperature, and then cBN grain rearrangement and further transformation of the interfacial layer are induced under increased pressure and temperature, making the nano-reactive layer formed in the previous step more dense and stable under high pressure, ultimately forming a robust ceramic-cBN transitional connection structure. Because these three steps provide a continuous process of interface coating, interface generation, and interface densification, the final ceramic-based PCBN composite material can effectively solve the defect of poor interface bonding. Attached Figure Description

[0015] Figure 1 This is a flowchart of a method for preparing ceramic-based PCBN composite material in one embodiment of the present invention. Detailed Implementation

[0016] To make the objectives, technical solutions, and advantages of this invention clearer, the invention will be further described in detail below with reference to the accompanying drawings and embodiments. It should be understood that the specific embodiments described herein are merely illustrative and not intended to limit the invention.

[0017] Please refer to Figure 1 This invention proposes a method for preparing ceramic-based PCBN composite materials, the steps of which include: S1. The ceramic matrix and titanium source are placed in an inert gas environment and ball-milled together. Then, cubic boron nitride powder and functional additives are added and heated to react, thereby obtaining a composite precursor. The functional additives include an amorphous carbon source.

[0018] The ceramic matrix comprises SiAlON powder and alumina micropowder. SiAlON powder is a multi-element solid solution ceramic material composed of nitrogen, oxygen, silicon, and aluminum, representing an industrial upgrade of the silicon nitride-based ceramic system. It possesses high strength, high toughness, high thermal stability, and good thermal shock resistance. In its crystal structure, silicon atoms can be partially replaced by aluminum atoms, and nitrogen atoms can be compensated by oxygen atoms, thus forming stable Si-N, Al-N, and Si-O bonds, enabling it to maintain excellent mechanical and thermal properties at high temperatures. In the ceramic matrix system of this invention, SiAlON serves as the basic ceramic phase, acting as a reinforcing framework. By constructing a three-phase ceramic network together with alumina and titanium sources, it promotes neck formation during the pre-sintering and high-temperature densification stages and exhibits good chemical stability and thermal compatibility with the cubic boron nitride interface, thereby providing a uniform surface energy environment for the deposition of the amorphous carbon reaction layer at the interface. The SiAlON powder used in all embodiments of the present invention is purchased separately from professional ceramic powder suppliers. Typical specifications include a purity greater than 99%, a particle size distribution between 0.5 and 2 micrometers, a morphology of spherical or irregularly broken, and good flowability and high specific surface area, making it suitable for thorough mixing with alumina and titanium sources during ball milling, thereby ensuring the uniformity of subsequent reactions and structural stability.

[0019] Step S1 includes: S1.1 Add SiAlON powder, alumina micro powder and titanium source powder to a ball mill jar in a mass ratio of (40-60):(10-20):(10-20), introduce argon or nitrogen to form an inert protective atmosphere, and use a planetary ball mill at 200-250 rpm for 4-6 hours to obtain mixed powder.

[0020] Step S1.1 involves prolonged ball milling of SiAlON powder, alumina micropowder, and titanium source powder under inert gas protection. This process ensures uniform mixing of the three ceramic phases under continuous shear, impact, and frictional energy, while simultaneously creating a slight coarsening and activation effect on the particle surface. The nitrogen-oxygen covalent network structure of SiAlON generates more exposed coordination sites during ball milling, which is beneficial for subsequent solid-phase diffusion. Alumina micropowder, with its high surface polarity, enhances the adsorption between ceramic particles, enabling rapid formation of ceramic necks during subsequent sintering. Titanium sources such as titanium nitride, titanium carbide, or titanium carbonitride, with their high hardness, act as micro-grinding media during ball milling, refining the overall particles and improving mixing uniformity. The formation of this uniform ceramic framework significantly reduces the diffusion resistance of subsequent interfacial reactions, improves the stability of the interfacial nano-reaction layer, and provides a homogeneous microscopic basis for the densification of the overall composite material, the construction of the interfacial structure, and the generation of defect networks.

[0021] Cubic boron nitride (cBN) powder is evenly spread in a ceramic boat, placed in a tube furnace with an inert atmosphere, and heated at 500-600°C for 0.5-1 hour with pure nitrogen or argon continuously introduced.

[0022] By thermally activating cubic boron nitride (cBN) powder at 500–600 °C under a pure nitrogen or argon atmosphere, slight atomic migration occurs on the surface of cBN grains, preferentially forming shallow point defects such as boron and nitrogen vacancies. These vacancy defects provide more migration pathways for subsequent interfacial chemical reactions, making it easier for amorphous carbon sources to deposit on the cBN surface during the pre-sintering stage. Furthermore, this promotes lattice rearrangement under high temperature and pressure conditions, thereby enhancing the tendency for subsequent twin boundaries and stacking fault structures to form. This method fully utilizes the thermal activation sensitivity of cBN, enabling the material to possess a controllable defect foundation in the precursor stage, which helps to construct a more stable interfacial structure and defect synergistic network.

[0023] Cubic boron nitride powder is placed in a low-pressure plasma reaction chamber, and nitrogen or argon gas is introduced at a radio frequency power of 50-150W to form a plasma state. The processing time is 5-15 minutes.

[0024] Surface activation of cubic boron nitride (cBN) powder using nitrogen plasma can rapidly introduce surface dangling bonds, dislocation sites, and shallow defect structures. These thermally stable electronic activation sites significantly enhance the interface's adsorption capacity for amorphous carbon precursors, allowing the carbon source to bind more uniformly and firmly to the cBN surface during subsequent heating stages. This results in a more continuous and controllable amorphous carbon nanolayer reaction layer. The plasma-induced active defects also serve as lattice slip initiation points during high-temperature, high-pressure sintering, promoting stacking faults and twinning in cBN, further strengthening the synergistic toughening effect of the composite material's interface structure.

[0025] The cubic boron nitride powder and ceramic grinding beads were placed together in a ball mill and ball-milled at a speed of 300-400 rpm for 10-30 minutes.

[0026] Mechanical etching of cubic boron nitride (cBN) powder using high-energy, short-time ball milling can create nanoscale roughness, microcracks, and mild lattice damage on the powder surface, significantly enhancing its surface free energy and interfacial reactivity. The roughened interface not only facilitates the capture and deposition of amorphous carbon sources on the cBN surface but also strengthens the intercalation between ceramic and cBN particles during the subsequent medium-temperature pre-sintering stage, significantly enhancing the mechanical locking ability between the interfaces. Furthermore, the shallow dislocations introduced by mechanical etching can become initiation points for lattice slip during the high-temperature, high-pressure stage, contributing to the induction of TB and SF defect structures and further improving the synergistic strength and toughness properties of the material.

[0027] S1.2 Add cubic boron nitride powder to a ball mill jar at 15-30 wt% of the mixed powder mass, maintain an inert protective atmosphere of argon or nitrogen, and continue ball milling for 1-2 hours at 150-180 rpm to obtain a composite mixed powder containing an ultrahard phase.

[0028] Step S1.2 involves adding pretreated cubic boron nitride (cBN) powder to the mixed ceramic matrix and continuing low-speed ball milling under an inert atmosphere. This helps to uniformly embed cBN into the ceramic framework composed of SiAlON, alumina, and titanium source. Since the pretreatment has created vacancy defects, surface roughness, or activation sites on the cBN surface, the ball milling process further enhances the interfacial contact between cBN and ceramic particles, resulting in a more stable multiphase mixed system. This structure can improve the deposition uniformity of the amorphous carbon reaction layer in the subsequent pre-sintering stage and preferentially propagate the interfacial reaction along the ceramic-cBN interface, laying the microscopic foundation for the subsequent formation of the AIL interfacial layer and defect structures such as twin boundaries and stacking faults.

[0029] S1.3 Add the functional additives to a ball mill jar at 3-10 wt% of the composite powder mass, and heat and stir at 80-120℃ to obtain a composite powder with an amorphous carbon precursor coated on the surface. The functional additives include an amorphous carbon source and an interface activator. The amorphous carbon source includes at least one of phenolic resin, polyimide, carbon black powder, glucose or sucrose solid carbon source. The interface activator includes at least one of lithium nitride, magnesium nitride, lithium borate, calcium fluoride, lithium hydroxide, lithium nitride, and lithium nitrate. The titanium source includes at least one of titanium nitride, titanium carbide, and titanium carbonitride.

[0030] Step S1.3 introduces amorphous carbon sources and interface activators under heating and stirring conditions, enabling these two types of functional additives to form a uniform adsorption layer on the surface of ceramic-cBN particles. Carbon sources such as phenolic resin, polyimide, carbon black, or glucose can soften or activate and coat the particle surface during low-temperature heating, providing a uniform carbon layer foundation for subsequent thermal decomposition. Interface activators such as lithium nitride, magnesium nitride, and lithium borate promote interfacial wetting, reduce diffusion barriers, and enhance the chemical activity of the ceramic-cBN interface. This step, through carbon source pre-coating and chemical activation, makes the amorphous carbon reaction layer and interface activation reaction more likely to occur in subsequent heat treatment stages, resulting in a denser interfacial structure, more complete reactions, and improved interfacial bonding strength and defect structure stability of the final composite material.

[0031] S1.4 Place the composite powder coated with amorphous carbon precursor in a tube furnace under an inert atmosphere and heat it at 200-300℃ for 1-2 hours to obtain the composite precursor.

[0032] Step S1.4 involves curing the composite powder coated with the amorphous carbon precursor at 200–300°C under an inert atmosphere. This process allows for initial cross-linking and curing of the carbon source, fixing the ceramic particles, cBN superhard phase, and interface activator within a stable precursor network. This curing process significantly improves the bonding strength of the carbon source at the ceramic-cBN interface, preventing interface inhomogeneity caused by carbon source migration during subsequent heating. Simultaneously, some of the interface activator undergoes initial reaction at this stage, enhancing the interfacial chemical activity in the subsequent pre-sintering stage. Ultimately, the stable precursor obtained in this step ensures the formation of a uniformly thick amorphous carbon nanolayer during the S2 thermal decomposition stage and supports the interface reconstruction and defect induction processes in subsequent high-temperature and high-pressure sintering, thereby improving the overall structural stability of the final ceramic-based PCBN composite material.

[0033] S2. The composite precursor is loaded into the forming mold and pre-sintered under pre-pressure to cause the amorphous carbon source to decompose and generate a nano-reaction layer between the cubic boron nitride powder and the ceramic matrix, thus obtaining the pre-sintered body.

[0034] Step S2 includes: S2.1. Fill the composite precursor into the silicon nitride mold and cold press it for 1 to 3 minutes with an initial forming pressure of 5 to 10 MPa to obtain the pressed blank.

[0035] Step S2.1 involves filling the composite precursor into a silicon nitride mold and applying initial cold pressing. Under external pressure, the large pores between the powder particles are gradually eliminated, forming a pressed green body with a basic shape and certain mechanical strength. Since SiAlON, alumina, and titanium source powders are all high-hardness ceramics, the pressing process allows for particle displacement, rolling, and local interlocking, thus forming an early mechanically locked structure. During cold pressing, the cubic boron nitride grains form multi-point contact with the ceramic framework, which is beneficial for increasing the effective contact area for subsequent interfacial chemical reactions. Furthermore, cold pressing allows the pre-coated amorphous carbon precursor to adhere more tightly to the ceramic phase and cBN surface, laying the foundation for continuous carbon layer deposition in the subsequent thermal decomposition stage. The technical effect of this step is to construct an initial green body with suitable density, sufficient interfacial contact, and uniform carbon source distribution, making the interfacial diffusion, carbon layer deposition, and grain connection in the subsequent pre-sintering process more stable and efficient.

[0036] S2.2 Place the pressed blank into a tube furnace and heat it at 700-900℃ for 10-60 minutes in an inert atmosphere of nitrogen, argon or helium to cause thermal decomposition of the amorphous carbon precursor, and form an amorphous carbon nano-reaction layer in situ at the interface of SiAlON, alumina and cubic boron nitride grains.

[0037] Step S2.2 Heating the pressed preform to 700–900°C under an inert atmosphere causes significant thermal decomposition of the amorphous carbon precursor, resulting in the formation of a mixture rich in sp2+.2 Bond-structured active carbon species are preferentially deposited at the interface between SiAlON, alumina, and cubic boron nitride particles, forming an amorphous carbon nanolayer with controllable thickness and good continuity. This interfacial carbon layer not only blocks the direct ceramic-cBN reaction and mitigates interfacial thermal mismatch, but also synergizes with the interfacial surfactant during subsequent high-temperature and high-pressure sintering, promoting diffusion, wetting, and lattice rearrangement at the interface. In particular, the nitrogen-oxygen coordination vacancies on the SiAlON and alumina surfaces, as well as the surface defects formed on the cBN surface after pretreatment, enable the active carbon to have stronger adsorption and binding effects at the interface, thereby improving the uniformity and adhesion stability of the carbon layer. The important technical effect of this step is to achieve selective deposition of the interface and self-assembly of the interfacial carbon layer, making it easier to generate favorable microstructures such as twin boundaries and stacking faults during subsequent sintering, significantly improving the toughness and interfacial stability of the composite material.

[0038] S2.3. The pressed blank with the amorphous carbon nano-reaction layer is kept at 700-900℃ and then subjected to a pressure of 5-10MPa for 5-15 minutes. After releasing the pressure, it is kept at 700-900℃ for 10-30 minutes to obtain the pre-sintered body.

[0039] Step S2.3 involves applying an external pressure of 5–10 MPa to the green body forming the nano-reaction layer while maintaining a temperature of 700–900°C. This further compacts the ceramic particles and cBN grains, enhancing the adhesion and continuity of the interfacial carbon layer under pressure. The pressure drives the active carbon species to penetrate further along interfacial defects, micropores, and grain boundaries, transforming the interfacial carbon layer from a simple deposition layer into a more dense composite reaction layer. Simultaneously, the ceramic phase undergoes slight solid-phase diffusion under pressure, resulting in a more stable necked structure between SiAlON and alumina, reducing porosity and improving overall densification. Continuing to maintain the temperature after releasing the pressure further homogenizes the microstructure, stabilizes the thickness of the interfacial carbon layer, and achieves a more balanced stress distribution between particles, ultimately forming a pre-sintered body with higher mechanical strength, stronger interfacial adhesion, and a more complete micro-reaction structure. The technical effect of this step is to complete the quality transformation of the interface structure from deposition to densification, which lays a key microscopic foundation for the ultra-hard phase connection, defect network construction and overall mechanical property improvement in the final high-pressure and high-temperature sintering stage.

[0040] S3. The pre-sintered body is first sintered under high temperature conditions, and then the ambient air pressure and temperature are increased to obtain ceramic matrix PCBN composite material.

[0041] Step S3 includes: S3.1. Load the pre-sintered body into the vacuum sintering furnace and evacuate to 10°C. -2 ~10 -3Pa, then the temperature is raised to 1300-1500℃ and held for 30-90 minutes to obtain the preform.

[0042] Step S3.1 involves high-temperature sintering of the pre-sintered body at 1300–1500℃ under high vacuum conditions. This effectively removes residual adsorbed gases, interfacial moisture, and low-temperature pyrolysis residues between particles, allowing the previously formed amorphous carbon nanoparticle reaction layer to stably adhere to the SiAlON, alumina, and cBN interfaces. Simultaneously, within this temperature range, significant solid-phase diffusion begins between SiAlON and alumina, forming a primary neck structure at the interface, significantly improving the strength and continuity of the ceramic framework. Titanium sources such as TiN, TiC, or Ti(C,N) exhibit excellent thermochemical stability at this stage, assisting in promoting interfacial diffusion and interphase bonding, resulting in a high density of the preform. This process also allows for further rearrangement of the amorphous carbon layer, maintaining its continuity at the interface and providing the necessary microscopic basis for the subsequent graphitization transformation and interface defect induction in the high-temperature, high-pressure stage. Therefore, the technical effect of this step is to construct a uniform and dense ceramic framework and stabilize the interfacial carbon layer, while eliminating gaseous impurities that affect high-pressure sintering, laying a crucial foundation for final microstructure refinement and phase boundary control.

[0043] S3.2. The precast body is placed into the ultra-high pressure cubic pressure chamber and heated to 1500-1700℃ under a pressure of 4-6 GPa, and kept at that temperature for 5-20 minutes.

[0044] Step S3.2, placing the preform in a cubic pressure chamber and performing high-temperature, high-pressure sintering under the coexistence of 4–6 GPa and 1500–1700 °C, is the core step in this invention for achieving structural reconstruction of the ceramic-based PCBN composite material. The ultra-high pressure stabilizes the cBN grains within the cubic structure region, and with the wetting assistance of the interfacial carbon layer, forms a strong interfacial bond with the ceramic framework. Simultaneously, the high temperature promotes the slip and rearrangement of the cBN lattice along specific directions, thereby inducing the formation of beneficial reinforcing structures such as stacking faults and twin boundaries. The amorphous carbon nanolayer reacts under high pressure to move towards high-density spp... 2 The directional graphitization of the structure enhances the thermochemical compatibility between the ceramic and superhard phases and absorbs interfacial stress. The combined participation of SiAlON, alumina, and titanium-based ceramics creates a three-phase synergistic support structure at the interface, improving interfacial mechanical stability and crack resistance. The technical effect of this step is to achieve cBN crystal structure optimization, interfacial carbon layer graphitization, and high-pressure densification of the ceramic framework, thereby constructing a composite microstructure with high strength, high toughness, and high heat resistance.

[0045] S3.3 Gradually release the pressure of 4-6 GPa and slowly cool it to room temperature at a rate of 5-20℃ / min to obtain ceramic matrix PCBN composite material.

[0046] Step S3.3, after the HPHT reaction is completed, gradually releasing the pressure by 4–6 GPa and slowly cooling at a rate of 5–20 °C / min effectively suppresses interfacial stress concentration caused by differences in thermal expansion coefficients, preventing thermal cracking of the ceramic phase, cBN grains, and interfacial carbon layer. The slow cooling process allows for sufficient relaxation of stress within the SiAlON, alumina, and cBN crystals, stabilizing and solidifying high-energy structures such as stacking faults and twin boundaries, preventing their destruction during rapid cooling. Furthermore, the interfacial graphitized carbon layer further achieves stress matching with the ceramic framework during slow cooling, thereby maintaining interfacial bonding strength and reducing the risk of phase boundary debonding. Ultimately, a ceramic-based PCBN composite material with intact microstructure, stable phase boundaries, and high strength and wear resistance is obtained.

[0047] This invention proposes a ceramic-based PCBN composite material, which is prepared by the method described above.

[0048] Example 1: First, SiAlON powder, alumina micro powder, and titanium nitride powder were weighed and added to a ball mill jar at a mass ratio of 50:20:20. Nitrogen gas was introduced to create an inert environment, and the mixture was continuously ball-milled at 230 rpm for 5 hours using a planetary ball mill to ensure uniform mixing of the three-phase ceramic base powders and the formation of activated contact interfaces. Then, 20 wt% of cubic boron nitride powder, pre-activated at 550℃ for 0.8 hours in an inert atmosphere, was added to the ball mill jar. The mixture was then ball-milled again at 160 rpm for 1.5 hours under an argon atmosphere to ensure uniform dispersion of the ultrahard phase within the ceramic system. Next, 8 wt% of phenolic resin was added as an amorphous carbon source, and an additional 2 wt% of lithium nitride was added as an interface activator. The mixture was heated and stirred at 100℃ for 40 minutes to uniformly coat the surfaces of SiAlON, alumina, and cubic boron nitride. The powder was then placed in a tube furnace filled with argon and heated at 250°C for 1.2 hours to allow the phenolic resin to undergo preliminary curing and be fixed into a stable carbon precursor, thus obtaining a composite precursor.

[0049] Next, the composite precursor was filled into a silicon nitride mold and cold-pressed at 7 MPa for 2 minutes to form a regular and dense pressed preform. The preform was then placed in a helium-protected tube furnace and heated at 800°C for 40 minutes to cause the phenolic resin to decompose and form an amorphous carbon nanolayer in situ at the interface of SiAlON, alumina, and cubic boron nitride. Subsequently, an external pressure of 8 MPa was applied again at the same temperature and held for 10 minutes to make the interfacial carbon layer more dense. The preform was then held at 800°C for another 20 minutes to form a structurally stable pre-sintered body.

[0050] Finally, the pre-sintered body is placed into a vacuum sintering furnace, and the furnace chamber is evacuated to 10°C. -3The material is heated to 1450℃ and held for 60 minutes to allow the ceramic phase to fully neck and complete the first densification. The preform is then placed in an ultra-high pressure cubic pressure chamber and held at 5 GPa and 1650℃ for 10 minutes to induce slip rearrangement of cubic boron nitride grains and induce TB and SF defect structures. At the same time, the interfacial carbon layer undergoes directional graphitization. The material is then slowly cooled to room temperature at a rate of 10℃ / min to finally obtain the ceramic matrix PCBN composite material.

[0051] Example 2: In this embodiment, SiAlON, alumina, and titanium carbide were mixed in a mass ratio of 45:15:20 and added to a ball mill jar. The mixture was ball-milled at 210 rpm for 6 hours under argon protection to form a homogeneous ceramic mixture. Cubic boron nitride powder, treated with 120W argon plasma for 10 minutes, was added at 25 wt% of the mixed powder mass and ball-milled at 160 rpm for 1 hour to uniformly embed the activated superhard phase into the ceramic framework. Then, 6 wt% of polyimide and 3 wt% of lithium nitrate were added, and the mixture was heated and stirred at 110°C for 50 minutes to uniformly coat the powder interface with polyimide. Subsequently, the mixture was cured at 280°C for 1 hour under nitrogen to obtain a composite powder with a carbon precursor coating. In the pre-sintering stage of this embodiment, the pressing process involves cold pressing at 6 MPa for 2 minutes to form the preform, followed by thermal decomposition at 850°C for 30 minutes to generate a carbon layer. A melamine-boric acid composite sol (0.5–2 wt%) is then sprayed onto the surface of the pressed preform with the amorphous carbon nanoparticle reaction layer, and the preform is held at 680–750°C for 5–10 minutes to allow the precursor to form a B–C–N-rich disordered transition phase at the interface, providing a slip surface for the subsequent high-pressure induced TB / SF structure. Then, a pressure of 7 MPa is applied at 850°C and held for 8 minutes, followed by a further 15 minutes to complete the pre-sintering. Finally, after 10… -2 It was held at 1380℃ for 40 minutes under vacuum conditions, then sintered at 4.5GPa and 1600℃ for 12 minutes, and then slowly cooled at a rate of 12℃ / min.

[0052] Example 3: In this example, SiAlON, alumina, and titanium carbonitride were mixed in a ratio of 55:15:15 and ball-milled at 240 rpm for 4 hours under nitrogen to generate a ceramic mixture. Cubic boron nitride was pretreated by short-time ball milling, in which cBN powder and ceramic grinding beads were placed together in a ball mill and subjected to high-energy mechanical treatment at a speed of about 350 rpm for about 20 minutes. Then, 18 wt% of the mixed powder was added, and ball milling continued at 170 rpm for 1 hour to form a composite powder. Glucose and calcium fluoride were added at 5 wt% of the composite powder, stirred at 90°C for 30 minutes, and then cured at 220°C for 1 hour to obtain the carbon precursor powder.

[0053] Subsequently, it was cold-pressed at 9 MPa for 1 minute to form a carbon layer, then thermally decomposed at 750℃ for 35 minutes to generate a carbon layer. After holding at 6 MPa for 10 minutes at 750℃, it was held for 20 minutes to form a pre-sintered body. Then, it was further... -3 A preform was obtained by holding it at 1500℃ for 80 minutes under vacuum (GPa). 0.2–1 wt% of a metal catalyst powder was added, and the preform was held at 900–1100℃ for 10–20 minutes to promote the local ordering of amorphous carbon at the interface, resulting in a short-range graphitization structure in the carbon layer. The preform was then held at 1700℃ for 8 minutes under 5.5 GPa, and cooled at a rate of 8℃ / min to obtain the final material. The metal catalyst powder included at least one of nickel-based metal catalyst powder, iron-based catalyst powder, and cobalt-based catalyst powder. The nickel-based catalyst powder included at least one of nickel powder, nickel-cobalt alloy powder, and nickel-molybdenum alloy powder. The iron-based catalyst powder included at least one of iron powder, iron-nickel alloy powder, and iron-cobalt alloy powder. The cobalt-based catalyst powder included at least one of cobalt powder, cobalt-nickel alloy powder, and cobalt-molybdenum alloy powder.

[0054] Comparative Example 1: The only difference between this comparative example and Example 1 is that in S1.3, only 8 wt% phenolic resin is used as the carbon source for the functional additives, and 2 wt% lithium nitride is not added. The remaining steps and proportions are exactly the same.

[0055] Comparative Example 2: The only difference between this comparative example and Example 1 is that no cubic boron nitride surface activation treatment was performed before S1.2; all other steps are the same as in Example 1.

[0056] Comparative Example 3: In this comparative example, the entire preparation process is consistent with that of Example 1. The only difference is that in step S1.3, no amorphous carbon source such as phenolic resin, polyimide, carbon black powder, glucose or sucrose solid carbon source is added, nor is any substance that can be cracked to generate a carbon layer used.

[0057] When characterizing the performance of the ceramic-based PCBN composite materials obtained in the examples and comparative examples, the surface hardness of the samples was first measured using the Vickers hardness test method. Each sample was processed into a flat test piece with a thickness of approximately 3–5 mm. An indentation test was performed using a Vickers hardness tester under a loading condition of HV 0.5. By applying a fixed load and holding it for approximately 10–15 seconds, a quadrilateral indentation was formed on the surface of the test piece. Subsequently, the diagonal length of the indentation was read using a microscope, and the Vickers hardness was calculated based on the geometric parameters of the indentation. To ensure data stability, each sample was tested 5 times, and the average value was taken as the final hardness result.

[0058] Fracture toughness testing was conducted using the indentation crack method. After polishing the sample surface to remove any obvious machining scratches, an indentation was formed by applying a Vickers indenter. In highly brittle materials, radial cracks would appear at the edge of the indentation. The fracture toughness value of the material could be calculated by measuring the length of the crack extending from the center of the indentation and substituting it into the relevant crack propagation formula. The same loading mode and indenter holding time were maintained throughout the testing process to ensure comparability of crack propagation behavior. Multiple measurements were performed on each sample to ensure data consistency.

[0059] The relative density test employs Archimedes' method. The sample is cut into a regular cubic or cylindrical structure, and its mass in air and its mass suspended in deionized water are measured using an analytical balance. The volume of the material is calculated based on the principle of buoyancy, and then the actual density is obtained by combining this with the mass in air. This actual density is then compared with the theoretical density to arrive at the relative density. During the test, the water temperature is maintained within the room temperature range, and air bubbles on the sample surface are eliminated to minimize density measurement errors.

[0060] The interfacial bond strength test employed a micro-tensile method. Sintered samples were cut into small strip specimens, with micro-clamps applied to both ends to prevent secondary damage to the loaded end faces. A micro-electromechanical tensile testing machine was then used to apply tension at a constant loading rate. The load at the moment of interfacial debonding was recorded, and the interfacial bond strength was calculated based on the stress area. Throughout the process, the samples were observed under a microscope to ensure that the fracture mode primarily reflected interfacial debonding behavior, ensuring that the test results accurately reflected the interfacial bonding performance between the ceramic phase and cubic boron nitride.

[0061] All the above tests were conducted at room temperature. The hardness, fracture toughness, density and interfacial bonding strength data obtained from the tests are the final performance indicators listed in Examples 1-3 and Comparative Examples 1-3. The experimental data are shown in Table 1.

[0062] Table 1: Example 1 introduces phenolic resin as an amorphous carbon source and incorporates 2wt% lithium nitride to enhance interfacial activity, enabling the formation of a continuous and dense carbon nanolayer in step S2. This effectively reduces the residual interfacial stress between SiAlON, alumina, and cBN, increasing the interfacial bonding strength to 220 MPa. The presence of the interfacial carbon layer improves load transfer efficiency, resulting in a transverse tensile strength of 1050 MPa. Simultaneously, the initial graphitization of the interfacial carbon layer significantly improves heat transfer efficiency, achieving a thermal conductivity of 38.5 W / m·K. The wear rate is reduced to 1.8 × 10⁻⁶. -6 mm 3 / N·m, exhibiting good wear resistance.

[0063] Example 2 introduces plasma activation before step S1.2 to generate more unsaturated B–N bonds on the cBN surface, improving interfacial reactivity. Then, after step S2.2, a melamine-boric acid sol is coated to form a BCN disordered layer, resulting in more slip surfaces at the interface. This structure more readily induces a TB / SF defect network under HPHT conditions, enhancing the dislocation slip capability between cBN grains. Consequently, the transverse fracture strength increases to 1120 MPa, and the interfacial bonding strength also rises to 245 MPa. The BCN transition layer improves the local thermal stability of the interface, increasing the thermal conductivity to 41.3 W / m·K and reducing the wear rate to 1.5 × 10⁻⁶. -6 mm 3 / N·m.

[0064] Example 3 employed short-time high-energy ball milling to activate cBN and introduced 0.2–1 wt% metal catalyst (Ni, Co, Fe-based). At 900–1100 °C, localized graphitization of amorphous carbon was induced, generating a short-range ordered sp² structure, giving the interfacial carbon layer graphite-like lamellar slip behavior. The metal catalyst significantly accelerated interfacial reconstruction, increasing the cBN-ceramic interfacial bonding strength to 260 MPa. The highly oriented carbon layer enhanced the chain-like thermally conductive network, achieving a thermal conductivity of 44.1 W / m·K. The wear rate was further reduced to 1.3 × 10 × 10⁻⁶. -6 mm 3 / N·m indicates that the material exhibits optimal interfacial strength and wear resistance.

[0065] In Comparative Example 1, the absence of lithium nitride resulted in discontinuous carbon layer formation at the interface, leading to intermittent carbon layers with poor adhesion after pyrolysis. Insufficient interfacial reaction between SiAlON and cBN resulted in a decrease in interfacial bonding strength to 155 MPa and transverse fracture strength to 880 MPa. Discontinuous heat conduction paths significantly reduced the thermal conductivity to 30.2 W / m·K, and increased the wear rate to 3.4 × 10⁻⁶. - 6 mm 3 / N·m indicates that the interface is fragile and the overall fatigue performance is reduced.

[0066] In Comparative Example 2, the cBN was not activated, and its surface lacked highly active unsaturated bonds, making it difficult for amorphous carbon to continuously deposit at the interface. This resulted in weak carbon layer adhesion and a decrease in interfacial bonding strength to 165 MPa. Unactivated cBN also had greater difficulty forming TB / SF defect structures during the HPHT stage, thus the material's hardness was only 40.5 GPa, and its wear rate was 3.1 × 10 × 10⁻⁶. -6 mm 3 / N·m. The results show that the activation treatment of cBN is a key step in the formation of the interfacial reaction layer.

[0067] In Comparative Example 3, no amorphous carbon source was added, resulting in the inability to form a carbon nanotube reaction layer in the S2 stage, leading to direct contact between the ceramic matrix and cBN. Due to the large difference in their thermal expansion coefficients and the high concentration of interfacial stress, the interfacial bonding strength decreased sharply to only 110 MPa, and the transverse fracture strength also dropped to 760 MPa. Simultaneously, due to the lack of slip channels provided by the interfacial carbon layer, it was difficult to induce TB / SF defect structures during the high-pressure sintering stage, resulting in an overall material thermal conductivity of only 25.3 W / m·K and a wear rate as high as 4.6 × 10⁻⁶. -6 mm 3 / N·m.

[0068] The above are merely preferred embodiments of the present invention and are not intended to limit the present invention. Any modifications, equivalent substitutions, and improvements made within the spirit and principles of the present invention should be included within the scope of protection of the present invention.

Claims

1. A method for preparing a ceramic-based PCBN composite material, characterized in that the steps include... include: S1. Place the ceramic matrix and titanium source in an inert gas environment, ball mill and mix them, then add cubic boron nitride powder and functional additives, and heat to react to obtain a composite precursor, wherein the functional additives include an amorphous carbon source. S2. The composite precursor is loaded into the forming mold and pre-pressure is applied for pre-sintering, so that the amorphous carbon source is decomposed and a nano-reaction layer is generated between the cubic boron nitride powder and the ceramic matrix to obtain the pre-sintered body. S3. The pre-sintered body is first sintered under high temperature conditions, and then the ambient air pressure and temperature are increased to obtain ceramic matrix PCBN composite material.

2. The method for preparing a ceramic-based PCBN composite material according to claim 1, characterized in that, The ceramic matrix includes SiAlON powder and alumina micro powder. Step S1 includes: S1.1 Add SiAlON powder, alumina micro powder and titanium source powder to a ball mill jar in a mass ratio of (40-60):(10-20):(10-20), introduce argon or nitrogen to form an inert protective atmosphere, and use a planetary ball mill at 200-250 rpm for 4-6 hours to obtain mixed powder. S1.2 Add cubic boron nitride powder to a ball mill jar at 15-30 wt% of the mixed powder mass, maintain an inert protective atmosphere of argon or nitrogen, and continue ball milling for 1-2 hours at 150-180 rpm to obtain a composite mixed powder containing an ultrahard phase. S1.3 Add the functional additives to the ball mill jar at 3-10 wt% of the composite powder mass, and heat and stir at 80-120℃ to obtain a composite powder with amorphous carbon precursor coated on the surface. S1.4 Place the composite powder coated with amorphous carbon precursor in a tube furnace under an inert atmosphere and heat it at 200-300℃ for 1-2 hours to obtain the composite precursor.

3. The method for preparing a ceramic-based PCBN composite material according to claim 2, characterized in that, The functional additives include amorphous carbon sources and interface activators. The amorphous carbon sources include at least one of phenolic resin, polyimide, carbon black powder, glucose or sucrose solid carbon sources. The interface activators include at least one of lithium nitride, magnesium nitride, lithium borate, calcium fluoride, lithium hydroxide, and lithium nitrate. The titanium sources include at least one of titanium nitride, titanium carbide, and titanium carbonitride.

4. The method for preparing a ceramic-based PCBN composite material according to claim 2, characterized in that, Step S1, preceding step S1.2, also includes: Cubic boron nitride powder is evenly spread in a ceramic boat, placed in a tube furnace with an inert atmosphere, and heated at 500-600°C for 0.5-1 hour with a continuous flow of pure nitrogen or argon.

5. The method for preparing a ceramic-based PCBN composite material according to claim 2, characterized in that, Step S1, preceding step S1.2, also includes: Cubic boron nitride powder is placed in a low-pressure plasma reaction chamber, and nitrogen or argon gas is introduced at a radio frequency power of 50-150W to form a plasma state. The processing time is 5-15 minutes.

6. The method for preparing a ceramic-based PCBN composite material according to claim 2, characterized in that, Step S1, preceding step S1.2, also includes: The cubic boron nitride powder and ceramic grinding beads were placed together in a ball mill and ball-milled at a speed of 300-400 rpm for 10-30 minutes.

7. The method for preparing a ceramic-based PCBN composite material according to claim 2, characterized in that, Step S2 includes: S2.

1. Fill the composite precursor into the silicon nitride mold and cold press it for 1 to 3 minutes with an initial forming pressure of 5 to 10 MPa to obtain the pressed blank. S2.2 Place the pressed blank into a tube furnace and heat it at 700-900℃ for 10-60 minutes in an inert atmosphere of nitrogen, argon or helium to cause thermal decomposition of the amorphous carbon precursor and in-situ deposition of an amorphous carbon nano-reaction layer at the interface of SiAlON, alumina and cubic boron nitride grains. S2.

3. The pressed blank with the amorphous carbon nano-reaction layer is kept at 700-900℃ and then subjected to a pressure of 5-10MPa for 5-15 minutes. After releasing the pressure, it is kept at 700-900℃ for 10-30 minutes to obtain the pre-sintered body.

8. The method for preparing a ceramic-based PCBN composite material according to claim 1, characterized in that, Step S3 includes: S3.

1. Load the pre-sintered body into the vacuum sintering furnace and evacuate to 10°C. -2 ~10 -3 Pa, then the temperature is raised to 1300-1500℃ and held for 30-90 minutes to obtain the preform; S3.

2. The precast body is placed into the ultra-high pressure cubic pressure chamber and heated to 1500-1700℃ under a pressure of 4-6 GPa, and kept at that temperature for 5-20 minutes. S3.3 Gradually release the pressure of 4-6 GPa and slowly cool it to room temperature at a rate of 5-20℃ / min to obtain ceramic matrix PCBN composite material.

9. A ceramic-based PCBN composite material, characterized in that, It is prepared by the method of preparing a ceramic matrix PCBN composite material as described in any one of claims 1-8.

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