High-toughness composite nitride ceramic blade, preparation method thereof and cutting tool

Abstract

This invention discloses a high-strength and high-toughness composite nitride ceramic blade, made of a sintered body containing a β-SiAlON phase and an equiaxed nitride phase. The sintered body meets the following technical conditions: the volume fraction of the equiaxed nitride phase is 30%~45%, the equiaxed nitride phase includes an α-SiAlON phase and / or an equiaxed Si3N4 phase, and the equiaxed nitride phase has grains with an aspect ratio not greater than 1.5; the solid solubility v value of the β-SiAlON phase satisfies 0.45≤v≤0.75, where v is the chemical formula Si. (6~V) Al V O V N (8~V) The variables include: the sintered body contains at least one rare earth element selected from La, Nd, Gd, Ho, Y, and Yb, which is introduced by rare earth oxides as sintering aids. The total content of all rare earth element oxides is 1.5~9.0 mol% based on the total molar amount of the sintered body; the average grain size of the sintered body is not greater than 0.8 μm; and the relative density of the sintered body is ≥99.5%. This ceramic blade exhibits excellent hardness and toughness at both room temperature and high temperature, reducing the risk of hot cracking and chipping.

Description

Technical Field

[0001] This invention belongs to the field of hard materials technology, and particularly relates to a high-strength and tough composite nitride ceramic cutting tool, its preparation method and cutting tool. Background Technology

[0002] SiAlON ceramics are solid solution materials composed of silicon, aluminum, oxygen, and nitrogen, formed by adding Al2O3 and AlN to Si3N4, causing the Si-N bonds to be replaced by Al-N and Al-O bonds during sintering. SiAlON ceramics possess unique comprehensive properties such as high-temperature stability, high strength, wear resistance, corrosion resistance, and thermal shock resistance, making them an ideal matrix material for manufacturing high-performance ceramic cutting tools. However, existing SiAlON ceramic cutting tools still face many technical challenges.

[0003] Traditional silon ceramic cutting tools exhibit low fracture toughness and insufficient resistance to thermal shock. While β-SiAlON phase possesses high fracture toughness, its hardness is relatively low; α-SiAlON, although possessing high hardness, suffers from low fracture toughness. A single crystal structure cannot simultaneously achieve both wear resistance and toughness. During high-speed cutting, the tool is prone to thermal cracking due to sudden temperature changes, leading to chipping and failure.

[0004] Therefore, there is an urgent need to provide a ceramic material with a novel reinforcement mechanism that can significantly improve toughness while maintaining high hardness. Summary of the Invention

[0005] The purpose of this invention is to provide a high-strength and tough composite nitride ceramic cutting tool, its preparation method and cutting tool, to solve the above-mentioned technical problems and improve the hardness and toughness of ceramic cutting tools.

[0006] To achieve this objective, the present invention adopts the following technical solution: In a first aspect, a high-strength and high-toughness composite nitride ceramic blade is provided, which is made of a sintered body containing a β-SiAlON phase and an equiaxed nitride phase, wherein the sintered body meets the following technical conditions: The volume fraction of the equiaxed nitride phase is 30%~45%, and the equiaxed nitride phase includes α-SiAlON phase and / or equiaxed Si3N4 phase. The equiaxed nitride phase is a grain with an aspect ratio of not more than 1.5. The solid solubility v of the β-SiAlON phase satisfies 0.45 ≤ v ≤ 0.75, where v is the chemical formula Si. (6~V) Al V O V N (8~V) Variables; The sintered body contains at least one rare earth element selected from La, Nd, Gd, Ho, Y, and Yb. The rare earth element is introduced by rare earth oxides as sintering aids, and the content of all rare earth element oxides is 1.5~9.0 mol% based on the total molar mass of the sintered body. The average grain size of the sintered body is no greater than 0.8 μm; The relative density of the sintered body is ≥99.5%.

[0007] Secondly, a method for preparing the above-mentioned high-strength and tough composite nitride ceramic cutting tool is provided, comprising the following steps: Raw material preparation: Weigh each raw material powder according to the above design ratio; Pretreatment: When the raw material contains α-Si3N4 powder, the α-Si3N4 powder is coated to form a nano-coating; when the raw material contains hard reinforcing particles, the hard reinforcing particles are pre-oxidized. Mixing and granulation: Mix the raw material powders, add the binder, and granulate; Molding: The granulated powder is pressed into shape to obtain a raw blank; Pre-firing: The green blank is pre-fired under a nitrogen atmosphere; Vibratory pressure sintering: The pre-fired green body is placed in a vibratory pressure sintering furnace and a nitrogen atmosphere is introduced; the temperature is raised and an initial pressure is applied; when the temperature reaches 1000~1200℃, the pressure is increased to 25~35MPa and kept constant, and a sinusoidal oscillating pressure with an amplitude of 2.5~5MPa and a frequency of 2~3Hz is superimposed on this pressure; the temperature is further raised to the sintering temperature of 1650~1800℃ and held for 30~90 minutes; after sintering, the oscillating pressure is stopped first, and a constant pressure is maintained until the temperature drops to 1000~1100℃, then the pressure is unloaded and cooled to room temperature; Post-processing: The sintered body is post-processed to obtain ceramic blades.

[0008] Thirdly, a cutting tool is provided, comprising the aforementioned high-strength and high-toughness composite nitride ceramic insert.

[0009] Compared with the prior art, the present invention has the following beneficial effects: This invention discloses a high-strength and high-toughness composite nitride ceramic cutting tool. Through a multiphase design of β-SiAlON phase (columnar crystals, providing toughening), and equiaxed nitride phase (equiaxed crystals, improving hardness and uniformity), combined with the grain boundary optimization effect of a specific rare earth element combination (at least one of La, Nd, Gd, Ho, Y, and Yb), the sintered body exhibits both high hardness and high toughness. The volume fraction of the equiaxed nitride phase is controlled at 30%~45%, the β-SiAlON solid solubility v value is controlled at 0.45~0.75, the average grain size is refined to below 0.8μm, and the relative density reaches above 99.5%. The fine-grained structure significantly enhances the material strength through the Hall-Page effect, and the high density effectively eliminates stress concentration, giving the ceramic cutting tool excellent hardness and toughness at both room temperature and high temperature, reducing the risk of hot cracking and chipping.

[0010] This invention discloses a method for preparing the aforementioned high-strength and tough composite nitride ceramic blade, comprising: raw material preparation, pretreatment, mixing and granulation, molding, pre-firing, oscillating pressure sintering, and post-treatment. In the oscillating pressure sintering step, a constant pressure of 25-35 MPa is applied at a temperature of 1000-1200°C, superimposed with a sinusoidal oscillating pressure of 2.5-5 MPa amplitude and 2-3 Hz frequency, followed by further heating to 1650-1800°C and holding for 30-90 minutes. The oscillating pressure, through dynamic driving force, promotes particle rearrangement, material diffusion, and plastic deformation, achieving a comprehensive effect of reducing sintering temperature, shortening sintering time, increasing density, and refining grain size, thereby improving material properties. This results in ceramic blades exhibiting excellent hardness and toughness at both room temperature and high temperature, reducing the risk of hot cracking and chipping.

[0011] The present invention provides a cutting tool comprising the above-mentioned high-strength and tough composite nitride ceramic insert, which exhibits excellent hardness and toughness at both room temperature and high temperature, reducing the risk of thermal cracking and chipping. Attached Figure Description

[0012] Figure 1 This is a flowchart illustrating the preparation process of the high-strength and tough composite nitride ceramic blade in this embodiment of the invention. Figure 2 This is a schematic diagram of the microstructure of the sintered body in an embodiment of the present invention. Detailed Implementation

[0013] The technical solution of the present invention will be further described below with reference to embodiments, but it is not limited thereto. Any modifications or equivalent substitutions to the technical solution of the present invention without departing from the spirit and scope of the present invention should be covered within the protection scope of the present invention. In the following embodiments, process equipment or devices not specifically specified are all conventional equipment or devices in the art. Unless otherwise specified, the raw materials used in the embodiments of the present invention are commercially available; unless otherwise specified, the technical means used in the embodiments of the present invention are all conventional means well known to those skilled in the art. Unless otherwise specified, the reagents, methods, and equipment used in the present invention are conventional reagents, methods, and equipment in the technical field.

[0014] This application provides a high-strength and tough composite nitride ceramic blade, which is made of a sintered body containing a β-SiAlON phase and an equiaxed nitride phase. The sintered body meets the following technical conditions: (1) The volume fraction of equiaxed nitride phase is 30~45%. The equiaxed nitride phase includes α-SiAlON phase and / or equiaxed Si3N4 phase. The equiaxed nitride phase is a grain with an aspect ratio of not more than 1.5. Equiaxed grains have uniform mechanical properties, which can reduce anisotropy and stress concentration inside the material and improve the toughness of the material.

[0015] (2) The solid solubility v of the β-SiAlON phase satisfies 0.45≤v≤0.75, where v is the chemical formula Si (6~V) Al V O V N (8~V) The variable value can be determined by X-ray diffraction. Within this range, SiAlON solid solution exhibits a stable crystal structure and excellent overall performance. Too low a v value leads to insufficient material strength, while too high a v value reduces the material's oxidation resistance and thermal stability. Rietveld-refined X-ray diffraction spectra can accurately determine the v value, ensuring the accuracy of the matrix composition.

[0016] (3) The sintered body contains at least one rare earth element among La, Nd, Gd, Ho, Y, and Yb. The rare earth element is introduced by rare earth oxides as sintering aids. The total content of all rare earth element oxides is 1.5~9.0 mol based on the total molar amount of the sintered body. La, Nd, Gd, Ho, Y, and Yb series rare earth element oxides are selected as sintering aids, with a total content of 1.5~9.0 mol. This series of rare earth oxides exhibits high ionic radius matching with Si and Al ions, enabling them to effectively dissolve in the SiAlON grain boundary phase, reduce grain boundary energy, and promote sintering densification. The introduction of La₂O₃, Nd₂O₃, Ho₂O₃, Gd₂O₃, Y₂O₃, and Yb₂O₃ further optimizes the high-temperature stability of the grain boundary phase. Synergistically with oscillating pressure, this promotes the formation and diffusion of the X(La, Nd, Gd, Ho, Y, Yb)-Si-Al-ON liquid phase, avoiding the high-temperature softening problem of the grain boundary phase caused by traditional single sintering aids. Preferably, the total content of La₂O₃, Nd₂O₃, Ho₂O₃, Gd₂O₃, Y₂O₃, and Yb₂O₃ accounts for 60-90 mol% of the total rare earth oxide content. This combination achieves a synergistic effect of sintering densification and grain boundary strengthening.

[0017] When the rare earth oxide content is below 1.5 mol% (lower limit): the amount of sintering liquid phase is insufficient, making densification difficult to complete. Furthermore, because the solid solubility of rare earth elements does not reach the critical solid solubility for a stable α-SiAlON phase, equiaxed crystals undergo a large-scale phase transformation into columnar β phase during sintering. This results in a sharp reduction in the volume fraction of the equiaxed nitride phase (below 30%). While the material retains some toughness, it loses high hardness and structural uniformity, exacerbating anisotropy and making it prone to thermal stress concentration and chipping during cutting. When the rare earth oxide content is above 9.0 mol% (upper limit): the solid solubility of rare earth elements reaches saturation. Excess rare earth cannot enter the crystal lattice but instead accumulates at grain boundaries, forming an excessively thick rare earth silicate glass phase. Due to the extremely low high-temperature strength of the amorphous glass phase, when the cutting temperature reaches above 1000℃, the excessively thick grain boundary phase softens or even locally liquefies, causing the grains to lose their bonding and bridging effect. The material's high-temperature strength, thermal shock resistance, and red hardness deteriorate sharply. During cutting, the cutting tool will undergo plastic deformation and adhesive wear caused by grain boundary softening, which will lead to rapid failure.

[0018] (4) The average grain size of the sintered body is no greater than 0.8 μm. Fine-grained structures can significantly improve the strength and hardness of materials through the Hall-Page effect.

[0019] (5) The relative density of the sintered body is ≥99.5%, with no obvious pores or cracks, ensuring a uniform internal structure and avoiding stress concentration. This can be determined by Archimedes' water displacement method.

[0020] In some embodiments, the sintered body further comprises hard reinforcing particles, the content of which is 2-18 mol% based on the total molar amount of the sintered body. The hard reinforcing particles are selected from at least one of Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, and HfN. This invention achieves a significant improvement in the overall performance of the SiAlON material by introducing a specific amount of hard reinforcing particles (such as Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, and HfN) into the SiAlON matrix. The reinforcement mechanism is mainly reflected in the following three aspects: (1) High hardness, high wear resistance, and high temperature performance: The selected hard reinforcing particles themselves have ultra-high hardness (such as Ti(C,N)≥24GPa) and high melting point. They are dispersed in the matrix and act as an anti-wear skeleton during the cutting process, directly bearing the cutting heat and friction stress, which significantly reduces the wear rate of the matrix and gives the cutting tool excellent red hardness and high temperature wear resistance.

[0021] (2) Thermal matching and residual stress toughening: The thermal expansion coefficient of the hard reinforcing particles is higher than that of the SiAlON matrix. During the sintering and cooling process, due to the difference in thermal expansion, a suitable amount of radial residual compressive stress is generated in the matrix around the particles. This microscopic residual compressive stress can hinder the opening and propagation of cracks. When the microcracks induced by the cutting process pass through the hard phase, the cracks will be deflected or passivated, thereby achieving particle toughening and making up for the lack of toughness of single-phase SiAlON ceramics.

[0022] (3) Interface bearing and load transfer: By pre-oxidizing the hard reinforcing particles, a uniform oxide layer of 5~50nm is formed on their surface. This transition layer effectively improves the wettability between the hard phase and the matrix, forming a strong chemical / metallurgical bonding interface. Under cutting stress, it can effectively transfer the load borne by the matrix to the high-modulus hard phase, avoiding chipping caused by stress concentration at the interface.

[0023] When the content of hard reinforcing particles is below 2.0 mol% (lower limit): the hard phase is too sparse in the matrix, failing to form a continuous and effective wear-resistant skeleton and load transfer network. At this point, the particles contribute little to improving the matrix hardness and wear resistance, and the toughening effect of residual compressive stress is not significant. The material is still prone to hot cracking and rapid wear under high-speed cutting. When the content of hard reinforcing particles is above 18.0 mol% (upper limit): the volume proportion of the hard phase in the matrix is ​​too high, making it extremely prone to severe agglomeration during ball milling and sintering. Agglomeration not only disrupts the continuity of the matrix but also exacerbates local thermal expansion mismatch, leading to huge local thermal stress concentrations when the cutting temperature changes abruptly, resulting in microcracks or even macrocracks. At the same time, excessively high hard phase addition hinders particle rearrangement and densification during sintering, leading to a decrease in relative density, a sharp increase in material brittleness, and a high risk of catastrophic chipping failure during cutting.

[0024] In some embodiments, the average particle size ratio of the hard reinforcing particles to the matrix ceramic particles is controlled between 1:1.5 and 1:3, wherein the matrix ceramic particles are ceramic grains in the sintered body other than the hard reinforcing particles. To avoid agglomeration of the reinforcing phase (i.e., hard reinforcing particles, hereinafter the same) and poor interfacial bonding, the size ratio of the hard reinforcing particles to the matrix ceramic particles is controlled between 1:1.5 and 1:3.

[0025] In some embodiments, the surface of the hard reinforcing particles undergoes a pre-oxidation treatment to form a uniform oxide layer with a thickness of 5-50 nm on the particle surface. This oxide layer can serve as an interfacial transition layer, improving the wettability and bonding strength between the reinforcing phase and the matrix, and enhancing load transfer efficiency. The pre-oxidation treatment conditions are: holding at 500-800°C in an air atmosphere for 0.5-2 hours.

[0026] In some embodiments, the grain size distribution of the sintered body satisfies D 90 / D 10 ≤2.5.

[0027] In some embodiments, the average coefficient of thermal expansion of the sintered body in the range of room temperature to 1200°C is not greater than 3.8 × 10⁻⁶. -6 / K. A low coefficient of thermal expansion reduces thermal stress caused by temperature changes during cutting, thus improving resistance to thermal shock.

[0028] In some embodiments, the room temperature thermal conductivity of the sintered body is not less than 15 W / (m·K), and the thermal conductivity at 500°C is not less than 80% of the room temperature thermal conductivity. Good thermal conductivity allows for rapid heat dissipation from the cutting zone, reducing tool temperature and minimizing thermal wear and cracking. The sintered body exhibits a fracture toughness ≥5.5 MPa·m¹ / ² and a Vickers hardness ≥18 GPa, combining both toughness and wear resistance.

[0029] In some embodiments, the room temperature three-point bending strength of the sintered body is not less than 900 MPa, and the three-point bending strength at 1000°C is not less than 750 MPa. This ensures that the tool has sufficient structural strength under both room temperature and high temperature cutting conditions.

[0030] In some embodiments, the total content of La2O3, Nd2O3, Y2O3, and Yb2O3 in the rare earth element oxides accounts for 60-90 mol of the total rare earth oxides.

[0031] In some embodiments, the grain boundary phase of the sintered body exhibits a composite structure of an amorphous phase encapsulating a nanocrystalline phase, wherein the nanocrystal size is 5-50 nm and the thickness of the amorphous phase does not exceed 2 nm. This can be confirmed by transmission electron microscopy (TEM), where the grain boundaries are mainly composed of a glassy phase consisting of Si, N, rare earth elements, and O.

[0032] In some embodiments, the sintered body further comprises Al2O3 nanoparticles distributed at the triple grain boundaries, wherein the content of the Al2O3 nanoparticles is 0.5-5 mol%, and the average particle size of the Al2O3 nanoparticles is 50-200 nm. The Al2O3 nanoparticles are preferentially distributed at the triple grain boundaries, which can achieve grain refinement by pinning grain boundaries and hindering grain growth, while simultaneously filling grain boundary voids and improving the material's density and thermal shock resistance. The distribution characteristics can be verified by scanning electron microscopy (SEM) combined with energy dispersive spectroscopy (EDS).

[0033] In some embodiments, the Al component in the sintered body is partially derived from AlON (aluminum oxynitride) or MgAl2O4 (spinel), wherein the content of AlON or MgAl2O4 accounts for 10-40 mol% of the total Al source. The purity of AlON is ≥99.5%, and the average particle size of spinel is ≤1.0 μm. The introduction of AlON and MgAl2O4 can further optimize the thermal conductivity and thermal shock resistance of the material, and they have good chemical stability and will not react harmfully with the matrix.

[0034] This application provides a high-strength and tough composite nitride ceramic cutting tool, which has the following performance characteristics: 1. Refined grain structure (average grain size ≤ 0.8 μm) and uniform distribution of reinforcing phase, with no obvious agglomeration; 2. Excellent wear resistance and chemical abrasion resistance, suitable for high-speed cutting of difficult-to-machine materials such as cast iron and heat-resistant alloys; 3. Room temperature strength (three-point bending strength ≥900MPa) and high temperature strength (1000℃ ≥750MPa), balancing red hardness and structural stability; 4. Excellent resistance to thermal shock, low coefficient of thermal expansion (≤3.8×10⁻⁶). -6 / K), high thermal conductivity (≥15W / (m·K) at room temperature), and can withstand sudden temperature changes during cutting; 5. High density (relative density ≥99.5%) and long service life, significantly improving tool life compared to traditional ceramic tools under high-speed dry cutting conditions.

[0035] This application also provides a method for preparing the above-mentioned high-strength and tough composite nitride ceramic blade, comprising the following steps: raw material preparation, pretreatment, mixing and granulation, molding, pre-firing, oscillating pressure sintering, and post-treatment. Specifically, as follows: S1. Raw material preparation: Weigh each raw material powder according to the design ratio.

[0036] In some embodiments, each raw material powder is precisely weighed according to the designed proportions, wherein the average particle size of α-Si3N4 powder is 0.2~1.0μm, purity ≥99.8%, and α phase content ≥95%; the average particle size of Al2O3 powder is 0.3~0.8μm, purity ≥99.9%; the average particle size of AlN powder is 0.8~1.2μm, oxygen content ≤0.5wt%; the average particle size of rare earth oxide powders (La2O3, Nd2O3, Ho2O3, Gd2O3, Y2O3, Yb2O3) is 0.3~1.5μm, purity ≥99.9%; the average particle size of hard reinforcing particle powders (Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, HfN) is 0.4~1.5μm, purity ≥99.5%; and the coefficient of variation of the loose density of all powders is ≤10%.

[0037] S2. Pretreatment: When the raw material contains α-Si3N4 powder, the α-Si3N4 powder is coated to form a nano-coating; when the raw material contains hard reinforcing particles, the hard reinforcing particles are pre-oxidized.

[0038] In some embodiments, the coating treatment is fluidized bed chemical vapor deposition coating treatment, and the thickness of the nano-coating formed on its surface is 5~50nm. The material of the nano-coating is Y2O3 or Al2O3, the coating coverage is ≥95%, and the coefficient of variation of coating thickness uniformity is ≤15%. The nano-coating can effectively prevent the oxidation of α-Si3N4 powder during mixing and sintering, while promoting the sintering activity of the powder. The pre-oxidation treatment conditions are: holding at 500~800℃ in an air atmosphere for 0.5~2 hours, heating rate of 5~10℃ / min, and naturally cooling to room temperature to form a surface oxide layer, which improves the interfacial bonding with the substrate.

[0039] S3. Mixing and granulation: Mix the raw material powders and add the binder, then granulate.

[0040] In some implementations, the mixing is ball milling.

[0041] In some embodiments, ball milling is carried out using a planetary ball mill with Si3N4 as the liner and grinding balls, anhydrous ethanol as the dispersion medium, a liquid-to-solid ratio of 1.5:1 to 2.5:1 (volume-to-mass ratio), a ball-to-material ratio of 15:1 to 25:1, a ball milling speed of 200 to 300 rpm, and a ball milling time of 24 to 72 hours, with a 30-minute stop every 12 hours to prevent powder overheating. After ball milling, 3 to 8 wt% of polyvinyl butyral (PVB) is added to the slurry as a binder, and after stirring evenly, spray drying is used for granulation. The inlet temperature is 180 to 220°C, and the outlet temperature is 80 to 100°C to obtain spherical granulated powder with good flowability. The average particle size of the granulated powder is 50 to 200 μm, and the flowability is ≤60 s / 50 g. A Si3N4 lined ball mill and grinding balls were used to avoid introducing impurities during the grinding process. Anhydrous ethanol was used as the dispersion medium, and the liquid-to-solid ratio and ball-to-material ratio were controlled. Through long-term ball milling, the uniform mixing and particle refinement of each component (including rare earth oxides of La, Nd, Gd, Ho, Y, and Yb, and hard reinforcing particles of Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, and HfN) were achieved. After adding PVB binder, spray drying granulation was used to obtain spherical granulated powder with good flowability, ensuring uniform filling of powder during the molding process and improving the density and uniformity of the green body.

[0042] In some embodiments, when the raw material contains AlON or MgAl2O4, the ball milling time is extended by 10 to 20 hours to ensure uniform dispersion of the Al component.

[0043] In some embodiments, granulation is carried out by spray drying granulation, with an inlet temperature of 180~220℃ and an outlet temperature of 80~100℃.

[0044] S4. Molding: Press the granulated powder into shape to obtain a raw blank.

[0045] In some embodiments, the pressing is cold pressing, with a pressing pressure of 100-200 MPa and a holding time of 30-60 seconds, balancing the green body density and molding efficiency. Granulated powder is uniformly loaded into a graphite mold, and a 0.1-0.3 mm thick layer of graphite emulsion is applied to the inner wall of the mold as a release agent. Cold pressing is performed under a pressure of 100-200 MPa for a holding time of 30-60 seconds to obtain a green body with a relative density of 55-65%, a dimensional accuracy error ≤ ±0.5%, and a coefficient of variation for the uniformity of internal pore distribution ≤ 12%.

[0046] S5. Pre-firing: The green blank is pre-firing under a nitrogen atmosphere. The pre-firing process is carried out under a nitrogen atmosphere, which removes the binder and forms preliminary green body strength, avoiding deformation or cracking of the green body during subsequent sintering.

[0047] In some embodiments, the pre-firing heating rate is 3~5℃ / min, the pre-firing temperature is 1000~1200℃, the holding time is 30~60 minutes, and the nitrogen flow rate is 50~100mL / min. The green blank is placed in a tube furnace and pre-firing is carried out in a nitrogen atmosphere (purity ≥99.9%), with the temperature increased to 1000~1200℃ at a rate of 3~5℃ / min and held for 30~60 minutes to remove the binder and form preliminary green body strength; during the pre-firing process, the nitrogen flow rate is controlled at 50~100mL / min, and the relative density of the pre-firing green body is 60~70%, with no obvious deformation or cracks.

[0048] S6. Oscillating pressure sintering: By controlling the oscillation pressure, the grain aspect ratio can be optimized, thus improving fracture toughness. Specifically, it includes the following steps: S61. Place the pre-fired billet in a vibrating pressure sintering furnace and introduce a nitrogen atmosphere.

[0049] Optionally, the oscillating pressure sintering furnace includes porous graphite electrodes, an infrared thermometer (temperature measurement accuracy ±5℃), a pressure sensor (measurement accuracy ±1%), a displacement sensor (monitoring shrinkage rate, accuracy ±0.1μm), and a temperature control system. Nitrogen gas (99.9% purity) is introduced into the sintering chamber as the sintering atmosphere, replacing the air in the chamber at least three times to ensure atmosphere purity. The nitrogen atmosphere is maintained throughout the sintering process to prevent material oxidation and component volatilization, ensuring the stability of the sintered body's chemical composition and inhibiting the high-temperature decomposition of silicon nitride.

[0050] S62. Increase the temperature and apply initial pressure; when the temperature reaches 1000~1200℃, increase the pressure to 25~35MPa and keep it constant, and then superimpose a sinusoidal oscillating pressure with an amplitude of 2.5~5MPa and a frequency of 2~3Hz.

[0051] In some embodiments, the heating rate is 10-15°C / min, and the initial pressure is 8-10 MPa; the cooling rate to room temperature is 8-10°C / min. The heating program is started at a heating rate of 10-15°C / min, while an initial pressure of 8-10 MPa is applied to pre-compress the powder.

[0052] S63. Continue heating to the sintering temperature of 1650~1800℃ and hold for 30~90 minutes. During this period, pressure feedback control is used to ensure stable oscillation pressure. In this temperature range, the transformation from α phase to β phase can be achieved with the assistance of oscillation pressure, and abnormal grain coarsening can be avoided.

[0053] In some implementations, the sintering temperature is 1650~1750℃ and the holding time is 30~60 minutes to ensure complete densification and phase transformation.

[0054] S64. After sintering, stop the oscillation pressure and maintain a constant pressure until the temperature drops to 1000~1100℃. Then unload the pressure and cool to room temperature. Specifically, it can be cooled to room temperature at a rate of 8~10℃ / min to avoid thermal stress cracks caused by sudden temperature changes.

[0055] Oscillating pressure sintering effectively promotes powder particle rearrangement, material diffusion, and plastic deformation through the continuous driving force generated by dynamic pressure. Compared with traditional sintering processes, it not only reduces the sintering temperature by 100-200℃ and shortens the sintering time, but also increases the density to over 99.5%. It promotes the uniform dispersion of reinforcing phases (containing Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, HfN) and rare earth elements such as La, Nd, Gd, Ho, Y, and Yb, resulting in a high-quality microstructure of "fine-grained matrix + uniformly dispersed reinforcing phase + optimized grain boundary structure" in the sintered body. At the same time, it improves fracture toughness and enhances the overall performance of the material by controlling the grain aspect ratio.

[0056] S7. Post-processing: The sintered body is post-processed to obtain ceramic blades.

[0057] In some embodiments, post-processing includes: removing the sintered body from the mold, grinding it with a diamond grinding wheel to remove the surface oxide layer and graphite contamination layer, and then performing chamfering to improve the dimensional accuracy and surface quality of the cutting tool.

[0058] In some implementations, the chamfer radius of the chamfering process is 0.1~0.3mm; the final ceramic insert has a dimensional accuracy error of ≤±0.02mm, a surface roughness Ra≤0.2μm, and a cutting edge sharpness of ≤5μm, ensuring cutting performance and surface quality during the cutting process.

[0059] This application also provides a cutting tool comprising the above-described high-strength and tough composite nitride ceramic insert.

[0060] Specifically, the prepared ceramic inserts can be fixed to a tool holder by welding or mechanical clamping to create cutting tools such as lathe tools, milling cutters, or boring tools. The optimized geometric parameters of the tool are: rake angle ~10°~5°, clearance angle 0°~15°, tip radius 0.2~1.2mm, chamfer width 0.1~0.3mm, and chamfer angle 10°~25°. This combination of parameters reduces cutting force and heat, and improves the tool's impact resistance and cutting stability. To further improve wear resistance, a multi-layer composite coating can be applied to the working surface of the insert, consisting of a transition layer (TiN or TiCN), a bonding layer (TiCNO), and a wear-resistant layer (Al2O3 or TiN) from the inside out. The coating adhesion force is ≥50N, ensuring that the coating does not peel off during high-speed cutting.

[0061] Examples and Comparative Examples

[0062] Raw material selection

[0063] The raw material powders used in the following examples and comparative examples, and their technical specifications, are as follows: α-Si3N4 powder: average particle size 0.5μm, purity ≥99.8%, α phase content ≥95%; Al2O3 powder: average particle size 0.4μm, purity ≥99.9%; AlN powder: average particle size 1.2μm, oxygen content ≤0.3wt%; Rare earth oxide powders (La2O3, Nd2O3, Ho2O3, Gd2O3, Y2O3, Yb2O3): average particle size 1.0~1.5μm, purity ≥99.9%; Hardened reinforced particulate powder (Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, HfN): average particle size 1.2μm, purity ≥99.5%; AlON powder: average particle size 0.8μm, purity ≥99.5%; Spinel (MgAl2O4) powder: average particle size 0.6μm, purity ≥99.5%.

[0064] In the following embodiments, unless otherwise specified, the preparation process is carried out according to the following uniform procedure: Pretreatment: α-Si3N4 powder was coated by fluidized bed chemical vapor deposition to form an Al2O3 nano-coating with a thickness of 20 nm and a coating coverage of ≥95% on its surface; the hard reinforcing particles were pre-oxidized by holding them at 600℃ in air for 1 hour at a heating rate of 8℃ / min and then naturally cooling them to room temperature to form a surface oxide layer.

[0065] Mixing and granulation: The raw material powders were placed in a planetary ball mill with Si3N4 as the liner and grinding balls. Anhydrous ethanol was added as the dispersion medium, with a liquid-to-solid ratio of 2:1 (volume-to-mass ratio) and a ball-to-material ratio of 20:1. The ball milling speed was 250 rpm, and the mixing was carried out for 48 hours, with a 30-minute stop every 12 hours to prevent the powder from overheating (when the raw materials contain AlON or spinel, the ball milling time is extended to 68 hours). After ball milling, 5 wt% polyvinyl butyral (PVB) was added to the slurry as a binder. After stirring evenly, the mixture was spray-dried and granulated at an inlet temperature of 200℃ and an outlet temperature of 90℃ to obtain spherical granulated powder with good flowability.

[0066] Molding: The granulated powder is evenly loaded into a graphite mold, and a 0.2 mm thick layer of graphite emulsion is coated on the inner wall of the mold as a release agent. Cold pressing is performed under a pressure of 150 MPa for 40 seconds to obtain a green blank.

[0067] Pre-firing: The green blank is placed in a tube furnace and pre-firing is carried out in a nitrogen atmosphere (purity ≥99.9%). The temperature is increased to 1100℃ at a rate of 4℃ / min and held for 45 minutes. The nitrogen flow rate is 80mL / min. This removes the binder and forms the initial green body strength.

[0068] Vibratory pressure sintering: The pre-fired green body is placed in a vibratory pressure sintering furnace, and nitrogen (99.9% purity) is introduced to replace the air in the chamber at least three times. A nitrogen atmosphere is maintained throughout the sintering process to prevent material decomposition. The heating program is started at a rate of 12℃ / min, while an initial pressure of 8MPa is applied to pre-compress the powder. When the temperature reaches 1100℃, the pressure is increased to 30MPa and kept constant. A sinusoidal oscillating pressure with an amplitude of 3MPa and a frequency of 2.5Hz is then applied. The temperature is further increased to the sintering temperature of 1700℃ and held for 45 minutes. After sintering, the oscillation pressure is stopped, and a constant pressure of 30MPa is maintained until the temperature drops to 1050℃. Then, the pressure is released, and the material is cooled to room temperature at a rate of 9℃ / min.

[0069] Post-processing: The sintered body is removed from the mold and ground with a diamond grinding wheel to remove the surface oxide layer and graphite contamination layer. Then, it is chamfered (chamfer radius 0.2mm). After grinding, the surface roughness Ra=0.15μm, the chamfer width of the cutting edge is 0.2mm, and the chamfer angle is 25°, finally obtaining the ceramic blade.

[0070] The raw material ratios (mol%) for each embodiment are shown in Table 1.

[0071] Table 1. Raw material ratios for the examples (mol%) Example 1 75.5 1.0 15.0 <![CDATA[4.5 La2O3]]> 4.0TaN – Example 2 78.5 1.5 15.0 <![CDATA[1.5 Nd2O3+1.5 Gd2O3]]> 2.0 Ti(C,N) – Example 3 76.0 1.5 16.0 <![CDATA[4.5 Ho2O3]]> 2.0 Ti(C,N) – Example 4 64.0 2.0 13.8 <![CDATA[2.0 La2O3+2.0 Nd2O3]]> 13.0 Ti(C,N) 3.2 AlON Example 5 74.8 2.5 7.45 <![CDATA[7.2 Gd2O3]]> 5.5 Ti(C,N) 2.55 Spinel Example 6 67.0 3.0 12.6 <![CDATA[6.0 La2O3]]> 6.0 Ti(C,N) 5.4 AlON Example 7 72.3 3.5 9.3 <![CDATA[8.0 Nd2O3]]> 6.9 NbN – Example 8 58.7 4.0 11.0 <![CDATA[7.0 Ho2O3]]> 14.0 TaN 5.3 Spinel Example 9 64.2 4.5 9.8 <![CDATA[3.5 Gd2O3+2.0 La2O3]]> 12.0 Ti(C,N) 4.0 AlON Example 10 70.8 2.0 10.7 <![CDATA[9.0 Nd2O3]]> 7.5NbC – Example 11 66.7 2.5 12.9 <![CDATA[3.5 La2O3]]> 2.5 Ti(C,N) +8.5 NbC 3.4 Spinel Example 12 68.0 1.5 12.75 <![CDATA[2.5 Ho2O3]]> 13.0 TaC 2.25 AlON Example 13 71.9 2.0 9.6 <![CDATA[5.5 Y2O3]]> 11.0 TiC – Example 14 59.7 2.3 13.5 <![CDATA[2.5 Yb2O3+1.5 Y2O3]]> 16.0 NbC 4.5 AlON Example 15 64.6 3.4 15.58 <![CDATA[3.0 Yb2O3]]> 6.5 Ti(C,N) + 3.5 TiC 3.42 Spinel Example 16 58.2 2.8 13.26 <![CDATA[1.5 La2O3+3 Yb2O3]]> 17.5 TaC 3.74 AlON Example 17 64.5 2.0 13.12 <![CDATA[4.5 Y2O3]]> 13.0 NbN 2.88 AlON Example 18 62.4 2.6 14.4 <![CDATA[3.0 Yb2O3]]> 12.5 TaN +1.5 TiN 3.6 Spinel Example 19 57.0 3.1 14.25 <![CDATA[2.0 La2O3+3.5 Y2O3]]> 15.5 HfN 4.65 AlON Example 20 69.0 1.5 15.0 <![CDATA[4.0 Nd2O3]]> 9.0 TiN +1.5 NbN – Example 21 60.5 2.5 14.45 <![CDATA[3.0 Gd2O3+2.5 Yb2O3]]> 12.5 TaN +2.0 HfN 2.55 Spinel Example 22 61.5 3.5 16.0 <![CDATA[6.5 Ho2O3]]> 3.0 NbN + 5.5 TiC 4.0 AlON Example 23 62.5 2.0 14.04 <![CDATA[2.5 La2O3+1.0 Y2O3]]> 12.0 HfN +2.0 NbC 3.96 Spinel Example 24 60.0 3.0 15.77 <![CDATA[2.0 Yb2O3+1.5 Nd2O3]]> 3.0 TiN + 11.5 TaN 3.23 AlON Note: For Examples 4, 5, 6, 8, 9, 11, 12, 14, 15, 16, 17, 18, 19, 21, 22, 23, and 24, the ball milling time was extended to 68 hours (48 hours + 20 hours) because the raw materials contained AlON or spinel; for the other examples that did not contain AlON or spinel, the ball milling time was 48 hours.

[0072] To verify the technical effect of the present invention, the following comparative examples were set up. The raw material ratio (mol%) and process parameters of each comparative example are shown in Table 2.

[0073] Table 2 Comparative Example Raw Material Ratio (mol%) and Process Parameters Comparative Example 1 83.0 1.0 15.0 <![CDATA[1.0 CeO2]]> 0 Hot pressing sintering 1750 Comparative Example 2 81.0 1.0 15.0 <![CDATA[3.0 Er2O3]]> 0 Hot pressing sintering 1750 Comparative Example 3 79.0 1.0 15.0 <![CDATA[5.0 Pr2O3]]> 0 Hot pressing sintering 1750 Comparative Example 4 81.0 1.0 15.0 <![CDATA[3.0 CeO2]]> 0 Oscillating pressure sintering 1750 Comparative Example 5 81.0 1.0 15.0 <![CDATA[3.0 Er2O3]]> 0 Oscillating pressure sintering 1750 Comparative Example 6 64.0 3.0 30.0 <![CDATA[3.0 CeO2]]> 0 Hot pressing sintering 1750 Comparative Example 7 79.0 2.0 14.0 <![CDATA[2.0 Pr2O3]]> 3.0 Ti(C,N) Oscillating pressure sintering 1750 Comparative Example 8 72.5 3.0 18.0 <![CDATA[2.5 Sm2O3]]> 4.0 NbC Oscillating pressure sintering 1750 Comparative Example 9 76.5 2.0 16.0 <![CDATA[2.5 Y2O3]]> 3.0 NbN (unoxidized) Oscillating pressure sintering 1750 Comparative Example 10 48.5 2.5 18.0 <![CDATA[3.0 Yb2O3]]> 28.0 TaN Oscillating pressure sintering 1750 Comparative Example 11 71.5 3.0 19.0 <![CDATA[2.0 La2O3+1.5 Y2O3]]> 3.0 HfN (unoxidized) Oscillating pressure sintering 1750 Comparative Example 12 78.5 3.3 15.0 <![CDATA[2.0 Nd2O3]]> 1.2 TiN Oscillating pressure sintering 1750 Comparative Example 13 47.0 2.5 17.0 <![CDATA[3.0 Gd2O3+0.5 Yb2O3]]> 15.0 NbN + 15.0 TaN Oscillating pressure sintering 1750 Comparative Example 14 73.7 3.5 18.5 <![CDATA[3.5 Ho2O3]]> 0.8 NbN Oscillating pressure sintering 1750 Comparative Example 15 73.0 2.0 18.0 <![CDATA[2.5 La2O3+1.0 Y2O3]]> 3.5 HfN (unoxidized) Oscillating pressure sintering 1750 Comparative Example 16 47.5 3.0 19.0 <![CDATA[2.0 Yb2O3+1.5 Nd2O3]]> 14.0 TiN+ 13.0 TaN Oscillating pressure sintering 1750 Note: Comparative Examples 1-3 and 6 used the traditional hot-pressing sintering process; Comparative Examples 4-5 and 7-16 used the oscillating pressure sintering process. Except for the differences noted in Table 2, the process parameters for each comparative example were the same as those for the examples. Specifically, the hard reinforcing particles in Comparative Examples 9, 11, and 15 were not pre-oxidized, while the pre-oxidation conditions for the hard reinforcing particles in the other comparative examples were the same as those in the examples. In Comparative Example 7, Pr₂O₃ is not within the rare earth element range defined by this invention; in Comparative Example 8, Sm₂O₃ is not within the rare earth element range defined by this invention; in Comparative Example 10, the amount of TaN added (28.0 mol%) exceeds the upper limit defined by this invention; in Comparative Example 12, the amount of TiN added (1.2 mol%) is lower than the lower limit defined by this invention; in Comparative Example 13, the total amount of NbN + TaN added (30.0 mol%) exceeds the upper limit defined by this invention; in Comparative Example 14, the amount of NbN added (0.8 mol%) is lower than the lower limit defined by this invention; and in Comparative Example 16, the total amount of TiN + TaN added (27.0 mol%) exceeds the upper limit defined by this invention.

[0074] Performance Testing and Result Analysis

[0075] Performance testing methods

[0076] (1) Microstructure analysis Grain morphology and reinforcing phase distribution were observed using scanning electron microscopy (SEM). The volume fraction of equiaxed nitride phase was determined by electron backscatter diffraction (EBSD). Equiaxed grains were defined as grains with an aspect ratio not greater than 1.5. Three different locations were selected at the center and edge of the sintered body for testing, and the average value was taken as the final result. Grain boundary structure and dislocation characteristics were analyzed using transmission electron microscopy (TEM). The sizes of no less than 500 grains were statistically analyzed using image analysis software, and cumulative distribution curves were plotted. The grain sizes corresponding to 10% and 90% cumulative distribution were recorded as D. 10 and D 90 The phase composition was analyzed by X-ray diffraction (XRD), and the solid solubility v value of the β-SiAlON phase was determined by Rietveld refined XRD spectra.

[0077] (2) Density measurement The relative density of the sintered body was determined using the Archimedes displacement method.

[0078] (3) Thermal performance testing The average coefficient of thermal expansion was determined using a thermal dilatometer in the range from room temperature to 1200°C. The thermal conductivity at room temperature and 500°C was determined using the laser scintillation method (JISR 1611).

[0079] (4) Mechanical property testing The three-point bending strength at room temperature and 1000℃ was determined using a universal testing machine with a three-point bending method (span 20 mm, loading rate 0.5 mm / min). Vickers hardness was determined using a Vickers hardness tester (load 30 kgf, holding pressure for 15 seconds). Fracture toughness was determined using the indentation method.

[0080] (5) Cutting performance test Ceramic inserts were machined to ISO SNGN120408 shape and mounted on a 100mm diameter face milling cutter. Dry cutting of FCD600 ductile cast iron (130mm × 700mm × 30mm) was performed. Cutting parameters: cutting speed 700m / min, feed rate 0.15mm / tooth, depth of cut 2.0mm. The number of workpieces cut was recorded based on the insert flank wear VB = 0.3mm or chipping.

[0081] Test Results The performance test results of each embodiment and comparative example are shown in Tables 3 and 4: Table 3. Test results of microstructure and thermal properties Example 1 31.2 0.45 0.68 2.2 3.45 99.6 17.5 85 Example 2 35.6 0.48 0.62 2.1 3.42 99.7 18.0 86 Example 3 29.8 0.51 0.71 2.3 3.48 99.6 17.2 84 Example 4 38.5 0.55 0.58 2.0 3.52 99.8 16.8 83 Example 5 41.2 0.61 0.55 2.0 3.58 99.7 15.5 82 Example 6 44.3 0.65 0.52 1.9 3.62 99.8 16.2 81 Example 7 39.8 0.68 0.60 2.1 3.65 99.7 16.0 80 Example 8 42.5 0.71 0.57 2.0 3.68 99.8 15.8 80 Example 9 45.2 0.75 0.53 1.9 3.72 99.7 15.5 81 Example 10 34.8 0.49 0.65 2.2 3.50 99.6 18.2 87 Example 11 37.2 0.53 0.60 2.0 3.55 99.8 17.8 85 Example 12 32.5 0.47 0.70 2.3 3.48 99.6 17.5 84 Example 13 36.5 0.50 0.63 2.1 3.46 99.8 18.5 88 Example 14 40.2 0.63 0.56 2.0 3.54 99.7 17.6 86 Example 15 38.8 0.66 0.54 1.9 3.57 99.8 17.3 85 Example 16 35.7 0.52 0.61 2.1 3.49 99.7 18.1 87 Example 17 37.6 0.54 0.60 2.1 3.51 99.7 17.9 86 Example 18 40.8 0.62 0.55 2.0 3.56 99.8 17.4 85 Example 19 43.5 0.67 0.53 1.9 3.63 99.8 17.0 84 Example 20 33.9 0.48 0.64 2.2 3.47 99.6 18.3 88 Example 21 42.1 0.64 0.56 2.0 3.59 99.7 16.9 83 Example 22 39.5 0.70 0.58 2.1 3.66 99.7 16.4 82 Example 23 38.2 0.58 0.59 2.0 3.55 99.8 17.7 85 Example 24 41.5 0.65 0.54 1.9 3.60 99.8 17.2 84 Comparative Example 1 24.2 0.29 1.25 3.1 3.18 98.2 13.2 75 Comparative Example 2 20.0 0.38 1.35 3.3 3.18 98.1 14.5 76 Comparative Example 3 23.3 0.39 1.20 3.2 3.20 98.3 19.0 74 Comparative Example 4 0 0.40 0.75 2.5 3.34 99.1 22.5 78 Comparative Example 5 19.1 0.41 0.82 2.6 3.20 99.0 12.1 77 Comparative Example 6 21.3 0.95 1.45 3.5 3.30 97.8 11.9 70 Comparative Example 7 27.5 0.43 0.88 2.7 3.32 99.2 14.8 79 Comparative Example 8 28.2 0.45 0.92 2.8 3.35 99.1 15.1 80 Comparative Example 9 30.5 0.52 0.85 2.7 3.40 98.8 15.3 81 Comparative Example 10 25.8 0.60 1.05 3.0 3.65 98.5 14.7 78 Comparative Example 11 29.2 0.55 0.90 2.8 3.45 98.7 15.0 80 Comparative Example 12 26.7 0.49 0.88 2.7 3.42 99.0 16.2 82 Comparative Example 13 24.3 0.63 1.10 3.2 3.70 98.3 14.2 77 Comparative Example 14 27.9 0.51 0.92 2.8 3.44 99.1 15.8 81 Comparative Example 15 28.6 0.57 0.95 2.9 3.48 98.6 15.2 79 Comparative Example 16 23.8 0.64 1.15 3.3 3.72 98.2 14.0 76

[0082] Table 4 Test results of mechanical and cutting properties Example 1 990 780 18.2 5.5 18 Normal wear and tear 0.28 Example 2 1020 800 18.5 5.7 22 Normal wear and tear 0.25 Example 3 1000 790 18.3 5.6 20 Normal wear and tear 0.27 Example 4 1050 820 18.8 5.8 25 Normal wear and tear 0.22 Example 5 1030 810 18.7 5.7 24 Normal wear and tear 0.23 Example 6 1010 800 18.6 5.6 23 Normal wear and tear 0.24 Example 7 1080 830 19.0 5.9 26 Normal wear and tear 0.21 Example 8 1060 820 18.9 5.8 25 Normal wear and tear 0.22 Example 9 1040 810 18.8 5.7 24 Normal wear and tear 0.23 Example 10 1100 840 19.2 6.0 28 Normal wear and tear 0.20 Example 11 1120 850 19.5 6.2 30 Normal wear and tear 0.18 Example 12 1020 805 18.5 5.7 22 Normal wear and tear 0.26 Example 13 1110 845 19.3 6.1 29 Normal wear and tear 0.19 Example 14 1130 855 19.6 6.3 31 Normal wear and tear 0.17 Example 15 1140 860 19.7 6.4 32 Normal wear and tear 0.16 Example 16 1125 842 19.4 6.2 30 Normal wear and tear 0.18 Example 17 1090 835 19.1 6.0 27 Normal wear and tear 0.2 Example 18 1115 848 19.4 6.2 29 Normal wear and tear 0.19 Example 19 1135 858 19.7 6.3 31 Normal wear and tear 0.17 Example 20 1070 825 18.9 5.9 26 Normal wear and tear 0.21 Example 21 1120 853 19.5 6.2 30 Normal wear and tear 0.18 Example 22 1085 832 19.0 5.9 27 Normal wear and tear 0.20 Example 23 1105 842 19.3 6.1 28 Normal wear and tear 0.19 Example 24 1128 856 19.6 6.3 30 Normal wear and tear 0.18 Comparative Example 1 780 650 17.0 4.2 7 Normal wear and tear 0.30 Comparative Example 2 700 660 17.2 4.3 10 Normal wear and tear 0.31 Comparative Example 3 700 650 17.1 4.2 8 Normal wear and tear 0.30 Comparative Example 4 750 580 18.0 4.5 3 chipping edge - Comparative Example 5 720 550 17.8 4.4 4 chipping edge - Comparative Example 6 720 520 16.5 4.0 3 rake face wear - Comparative Example 7 830 630 17.5 4.6 9 Normal wear and tear 0.30 Comparative Example 8 830 620 17.4 4.5 8 Normal wear and tear 0.32 Comparative Example 9 850 640 17.3 4.4 8 Normal wear and tear 0.30 Comparative Example 10 780 570 17.6 4.3 5 chipping edge - Comparative Example 11 820 620 17.4 4.5 7 Normal wear and tear 0.30 Comparative Example 12 860 650 17.7 4.6 9 Normal wear and tear 0.31 Comparative Example 13 760 560 17.2 5.0 4 chipping edge - Comparative Example 14 840 630 17.5 4.5 8 Normal wear and tear 0.36 Comparative Example 15 810 610 17.3 4.4 6 Normal wear and tear 0.30 Comparative Example 16 750 550 17.1 4.2 4 chipping edge -

[0083] Results Analysis The test results show that the composite nitride ceramic cutting tool of the present invention is significantly superior to the comparative example in terms of microstructure, thermal properties, mechanical properties and cutting performance.

[0084] (1) Microstructure The average grain size of the sample in the examples was 0.52~0.71 μm, much smaller than that of the comparative examples (0.75~1.45 μm), and the grain size distribution was more uniform (D). 90 / D 10 The volume fraction of the equiaxed nitride phase is 29.8%~45.2%, and the relative density is ≥99.5%. This is due to the dynamic densification effect of oscillating pressure sintering and the synergistic grain refinement effect of rare earth oxides such as La, Nd, Gd, Ho, Y, and Yb. The reinforcing phase is uniformly dispersed in the matrix without obvious agglomeration, and the interface bonding is tight.

[0085] In contrast, Comparative Examples 1-3, which employed traditional hot-pressing sintering, exhibited coarse grains (1.20-1.35 μm) and insufficient density (98.1%-98.3%). Comparative Examples 4-5, due to inappropriate selection of rare earth elements (CeO2, Er2O3), showed poor microstructural stability and significantly low volume fraction of equiaxed nitride phases (the α value in Comparative Example 4 was even 0%, indicating that the α-SiAlON phase could not be effectively formed). Comparative Examples 10, 13, and 16, due to excessive addition of hard phase (28-30 mol%), showed agglomeration and abnormal grain growth, with an average grain size exceeding 1.0 μm. Comparative Examples 9, 11, and 15, due to the lack of pre-oxidation treatment of the hard phase, showed poor interfacial bonding and microcracks at the interface.

[0086] (2) Thermal performance The thermal conductivity of the sample in the examples was 15.5~18.5 W / (m·K) at room temperature, the thermal conductivity retention rate at 500℃ was ≥80%, and the coefficient of thermal expansion was 3.42~3.72×10⁻⁶. -6 / K. The low expansion and high thermal conductivity give the material good resistance to thermal shock.

[0087] Comparative Examples 1-3 had thermal conductivity below 14.5 W / (m·K), and their thermal conductivity retention rate at 500℃ was only 74%-76%. Comparative Example 4, although using oscillating pressure sintering and having a relatively high thermal conductivity (22.5 W / (m·K)), suffered from insufficient high-temperature strength due to improper selection of CeO2 rare earth elements, resulting in chipping during cutting. Comparative Example 6, due to excessive AlN content leading to a v-value exceeding the range of this invention (0.95), had a thermal conductivity of only 11.9 W / (m·K), significantly reduced high-temperature stability, and its thermal conductivity retention rate at 500℃ was only 70%. Comparative Examples 10, 13, and 16, due to excessive addition of hard phase, had thermal conductivity reduced to 14.0-14.7 W / (m·K), and their coefficient of thermal expansion increased to 3.65-3.72 × 10⁻⁶. -6 / K.

[0088] (3) Mechanical properties The three-point flexural strength of the sample in the example was 980~1150 MPa at room temperature, 780~860 MPa at 1000℃, and the fracture toughness was 5.5~6.4 MPa·m. 1 / 2 All were significantly higher than the control group.

[0089] Comparative Examples 1-3, due to their coarse grains (1.20-1.35 μm) and insufficient density (98.1%-98.3%), exhibited room temperature strengths of only 700-780 MPa and fracture toughness of only 4.2-4.3 MPa·m. 1 / 2Comparative Examples 4 and 5, due to improper selection of rare earth elements (CeO2, Er2O3), exhibited a severe decrease in high-temperature strength to 550-580 MPa, resulting in chipping during cutting. Comparative Example 6, with a v value exceeding the range of this invention (0.95), showed increased material brittleness, with a room temperature strength of only 720 MPa and a high-temperature strength of only 520 MPa. Comparative Examples 7 and 8, due to improper selection of rare earth elements (Pr2O3, Sm2O3), although the amount of hard phase added was within the range of this invention, still achieved a strength of only 830 MPa and a toughness of 4.5-4.6 MPa·m. 1 / 2 Comparative Examples 9, 11, and 15, due to the lack of pre-oxidation treatment of the hard phase, exhibited poor interfacial bonding, resulting in a toughness of only 4.4–4.5 MPa·m. 1 / 2 Comparative Examples 10, 13, and 16 had increased material brittleness due to the addition of hard phase exceeding the upper limit of this invention (28-30 mol%), resulting in chipping during fracture and making it impossible to obtain effective VB data.

[0090] (4) Cutting performance The number of workpieces cut in the example samples ranged from 18 to 32, all exhibiting normal wear with a flank wear VB ≤ 0.28 mm. Example 15 showed the best overall performance (32 pieces, VB = 0.16 mm). It used Yb₂O₃ as a sintering aid, added a Ti(C,N) + TiC composite reinforcing phase, and introduced spinel as a partial Al source. Through an optimized oscillating pressure sintering process, it obtained a uniform fine-grained structure (average grain size 0.54 μm), high density (99.8%), and excellent mechanical properties (room temperature strength 1140 MPa, fracture toughness 6.4 MPa·m). 1 / 2 Example 19 uses La2O3+Y2O3 as a sintering aid and adds HfN hard reinforcing particles. It cuts 31 workpieces and its overall performance is second only to Example 15.

[0091] The comparative examples involved cutting only 3 to 10 workpieces. Comparative Examples 1 to 3 reached the failure criterion of VB=0.3mm after cutting 7 to 10 workpieces; Comparative Examples 4 to 5, 10, 13, and 16 experienced chipping and premature failure after cutting 3 to 5 workpieces; Comparative Example 6 exhibited abnormal rake face wear after cutting 3 workpieces; Comparative Examples 7 to 9, 11 to 12, and 14 to 15 reached the failure criterion of VB=0.30 to 0.36mm after cutting 6 to 9 workpieces. These results fully demonstrate that the cutting performance and service life of the ceramic inserts of this invention are significantly superior to existing products.

[0092] The above results demonstrate that the present invention, through reasonable material composition design (introducing rare earth oxides of La, Nd, Gd, Ho, Y, and Yb, and hard reinforcing particles of Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, and HfN) and innovative oscillating pressure sintering process, has successfully prepared composite nitride ceramic cutting tools with ultra-high wear resistance, high strength and toughness, and good thermal shock resistance, thus verifying the scientific nature and effectiveness of the technical solution of the present invention.

[0093] Selectable range of process parameters

[0094] Based on the above embodiments, the parameters of each process step of the present invention can be adjusted within the following ranges: (1) Raw material preparation The average particle size of α-Si3N4 powder can be 0.2~1.0μm, with a purity ≥99.8% and an α phase content ≥95%; the average particle size of Al2O3 powder can be 0.3~0.8μm, with a purity ≥99.9%; the average particle size of AlN powder can be 0.8~1.2μm, with an oxygen content ≤0.5wt%; the average particle size of rare earth oxide powder can be 0.3~1.5μm, with a purity ≥99.9%; the average particle size of hard reinforcing particles can be 0.4~1.5μm, with a purity ≥99.5%; and the coefficient of variation of the loose density of all powders is ≤10%.

[0095] (2) Pretreatment The coating of α-Si3N4 powder can be performed using fluidized bed chemical vapor deposition, with a coating thickness of 5-50 nm, using Y2O3 or Al2O3 as the coating material, and a coating coverage of ≥95%. The pre-oxidation treatment of the hard reinforcing particles can be carried out by holding them at 500-800℃ in air for 0.5-2 hours at a heating rate of 5-10℃ / min, followed by natural cooling to room temperature to form a surface oxide layer.

[0096] (3) Mixing and granulation Ball milling can be carried out in a planetary ball mill with Si3N4 as the liner and grinding balls, using anhydrous ethanol as the dispersion medium. The liquid-to-solid ratio is 1.5:1 to 2.5:1 (volume-to-mass ratio), the ball-to-material ratio is 15:1 to 25:1, the ball milling speed is 200 to 300 rpm, and the ball milling time is 24 to 72 hours, with a 30-minute stop every 12 hours to prevent powder overheating (when the raw material contains AlON or spinel, the ball milling time is extended by 10 to 20 hours). The binder can be polyvinyl butyral (PVB), with an addition amount of 3 to 8 wt%. The inlet temperature of spray drying granulation can be 180 to 220℃, and the outlet temperature can be 80 to 100℃. The average particle size of the obtained granulated powder is 50 to 200 μm, and the flowability is ≤60 s / 50 g.

[0097] (4) Molding The cold pressing pressure can be 100~200MPa, the holding time can be 30~60 seconds, and the relative density of the green blank is 55~65%.

[0098] (5) Pre-firing Pre-firing can be carried out under a nitrogen atmosphere (purity ≥99.9%), with a heating rate of 3~5℃ / min, a pre-firing temperature of 1000~1200℃, a holding time of 30~60 minutes, a nitrogen flow rate of 50~100mL / min, and a relative density of 60~70% for the pre-firing green body.

[0099] (6) Oscillating pressure sintering The heating rate can be 10~15℃ / min, and the initial pressure can be 8~10MPa. The pressurization temperature is 1000~1200℃, the constant pressure is 25~35MPa, the oscillation pressure amplitude is 2.5~5MPa, the frequency is 2~3Hz, and the waveform is a sine wave. The sintering temperature is 1650~1800℃, preferably 1650~1750℃; the holding time is 30~90 minutes, preferably 30~60 minutes. During the cooling stage, the constant pressure is first maintained to 1000~1100℃, and then the pressure is unloaded at a cooling rate of 8~10℃ / min.

[0100] (7) Post-processing The surface oxide layer and graphite contamination layer can be removed by grinding with a diamond wheel, followed by chamfering with a chamfer radius of 0.1~0.3mm. The final finished blade has a dimensional accuracy error of ≤±0.02mm, a surface roughness Ra≤0.2μm, and a cutting edge sharpness of ≤5μm.

[0101] The specific values ​​within the above-mentioned process parameter range can be adjusted by those skilled in the art according to actual needs and in conjunction with the teachings of the embodiments, all of which can achieve the technical effects of the present invention.

[0102] The above-described embodiments are only used to illustrate the technical solutions of the present invention, and are not intended to limit it. Although the present invention 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. Such 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 the present invention.

Claims

1. A high-strength and high-toughness composite nitride ceramic blade, characterized in that, It is made of a sintered body containing a β-SiAlON phase and an equiaxed nitride phase, and the sintered body meets the following technical conditions: The volume fraction of the equiaxed nitride phase is 30%~45%, and the equiaxed nitride phase includes α-SiAlON phase and / or equiaxed Si3N4 phase. The equiaxed nitride phase is a grain with an aspect ratio of not more than 1.

5. The solid solubility v of the β-SiAlON phase satisfies 0.45 ≤ v ≤ 0.75, where v is the chemical formula Si. (6~V) Al V O V N (8~V) Variables; The sintered body contains at least one rare earth element selected from La, Nd, Gd, Ho, Y, and Yb. The rare earth element is introduced by rare earth oxides as sintering aids, and the content of all rare earth element oxides is 1.5~9.0 mol% based on the total molar mass of the sintered body. The average grain size of the sintered body is no greater than 0.8 μm; The relative density of the sintered body is ≥99.5%.

2. The high-strength and high-toughness composite nitride ceramic blade according to claim 1, characterized in that, The sintered body further comprises hard reinforcing particles, and the content of the hard reinforcing particles is 2-18 mol% based on the total molar amount of the sintered body. The hard reinforcing particles are selected from at least one of Ti(C,N), NbC, TaC, TiC, NbN, TaN, TiN, and HfN.

3. The high-strength and high-toughness composite nitride ceramic blade according to claim 2, characterized in that, The surface of the hardened reinforcing particles is pre-oxidized to form a uniform oxide layer with a thickness of 5-50 nm on the particle surface.

4. The high-strength and high-toughness composite nitride ceramic blade according to claim 1, characterized in that, In the rare earth element oxides, the total content of La2O3, Nd2O3, Y2O3, and Yb2O3 accounts for 60-90 mol of the total rare earth oxides.

5. The high-strength and high-toughness composite nitride ceramic blade according to claim 1, characterized in that, In the sintered body, the interface between the β-SiAlON phase and the equiaxed nitride phase crystals presents a composite structure in which an amorphous phase encapsulates a nanocrystalline phase, wherein the nanocrystalline size is 5~50nm and the thickness of the amorphous phase does not exceed 2nm.

6. A method for preparing the high-strength and high-toughness composite nitride ceramic blade according to any one of claims 1 to 5, characterized in that, Includes the following steps: Raw material preparation: Weigh each raw material powder according to the design ratio according to any one of claims 1 to 5; Pretreatment: When the raw material contains α-Si3N4 powder, the α-Si3N4 powder is coated to form a nano-coating; when the raw material contains hard reinforcing particles, the hard reinforcing particles are pre-oxidized. Mixing and granulation: Mix the raw material powders, add the binder, and granulate; Molding: The granulated powder is pressed into shape to obtain a raw blank; Pre-firing: The green blank is pre-fired under a nitrogen atmosphere; Vibratory pressure sintering: The pre-fired green body is placed in a vibratory pressure sintering furnace and a nitrogen atmosphere is introduced. Heat the material and apply initial pressure. When the temperature reaches 1000~1200℃, increase the pressure to 25~35MPa and keep it constant. On this basis, superimpose a sinusoidal oscillating pressure with an amplitude of 2.5~5MPa and a frequency of 2~3Hz. Continue to heat the material to the sintering temperature of 1650~1800℃ and hold it for 30~90 minutes. After sintering, stop the oscillating pressure and maintain a constant pressure until the temperature drops to 1000~1100℃. Then, unload the pressure and cool to room temperature. Post-processing: The sintered body is post-processed to obtain ceramic blades.

7. The preparation method according to claim 6, characterized in that, The coating treatment is a fluidized bed chemical vapor deposition coating treatment, the thickness of the nano-coating is 5~50nm, and the material of the nano-coating is Y2O3 or Al2O3; the pre-oxidation treatment conditions are: heat preservation in air atmosphere at 500~800℃ for 0.5~2 hours.

8. The preparation method according to claim 6, characterized in that, The sintering temperature is 1650~1750℃, and the holding time is 30~60 minutes.

9. The preparation method according to claim 6, characterized in that, The heating rate is 10~15℃ / min, the initial pressure is 8~10MPa; the cooling rate to room temperature is 8~10℃ / min.

10. A cutting tool, characterized in that, The high-strength and high-toughness composite nitride ceramic blade includes any one of claims 1 to 5.

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