High red hardness ti(c,n)-based cermets tool material, its preparation method and tool

Abstract

The present application relates to other hard alloy manufacturing and superhard tool material preparation technical field, the present application relates to superhard tool material and its preparation technical field, specifically relates to a kind of high red hardness Ti (C, N) based cermet tool material and its preparation method and tool.The preparation method includes powder preparation, wet ball milling, drying granulation, compact, vacuum degreasing, vacuum sintering and controlled cooling, by making Ti (C, N) particle surface layer occurs selective dissolution and in subsequent cooling process at Ti (C, N) core periphery ring layer is formed to form ring core structure.The scheme can improve the wetting state and interface metallurgical bonding between Ti (C, N) ceramic phase and high specific gravity tungsten alloy binder phase, improve the high temperature hardness, oxidation resistance, transverse rupture strength and cutting life of material, suitable for high-speed cutting, dry cutting and difficult-to-cut material cutting field.

Description

Technical Field

[0001] This invention relates to the fields of other cemented carbide manufacturing and superhard tool material preparation technology, specifically to a high red hardness Ti(C,N)-based cermet tool material and its preparation method and tool. Background Technology

[0002] Ti(C,N)-based cermet tools, due to their high hardness, good wear resistance, and low coefficient of friction, have been widely used in the finishing and semi-finishing of steel, cast iron, and some high-temperature alloys. Existing Ti(C,N)-based cermets typically use Ni or Co as the metallic binder phase. However, with the increasing demands for cutting speed, machining accuracy, and the ability to handle difficult-to-machine materials in the manufacturing industry, traditional Ti(C,N)-Ni or Ti(C,N)-Co cermets are gradually revealing insufficient high-temperature performance under high-temperature cutting conditions. This is mainly manifested in the binder phase's tendency to soften, oxidize, and creep at high temperatures, leading to a decrease in tool red hardness, reduced wear resistance, and shortened tool life.

[0003] High-density tungsten alloys possess high density, high high-temperature strength, low coefficient of thermal expansion, and good creep resistance, theoretically making them suitable as high-temperature stable binder phases in Ti(C,N)-based cermets. However, directly introducing high-density tungsten alloys into the Ti(C,N) system introduces new technical challenges. Firstly, the high-melting-point tungsten alloy melt exhibits poor wettability with Ti(C,N) ceramic particles, making it difficult to fully fill interparticle gaps during sintering, easily leading to porosity defects and affecting the material's density and mechanical properties. Secondly, if an effective transition interface is lacking between the tungsten alloy and Ti(C,N), the bonding between them is weak, making them prone to interface debonding and early failure under cutting stress. Furthermore, in systems with high tungsten content, traditional sintering processes struggle to stably induce the formation of a complete and uniform ring-core structure, resulting in a trade-off between material strength and toughness.

[0004] Therefore, how to achieve good wetting and densification sintering between Ti(C,N) ceramic phase and high-density tungsten alloy binder phase under high tungsten content conditions, and how to stably construct a ring-core structure that is conducive to improving interfacial bonding strength and high-temperature service performance, remains a technical problem that urgently needs to be solved in this field. Summary of the Invention

[0005] The purpose of this invention is to provide a high red hardness Ti(C,N)-based cermet cutting tool material and its preparation method and cutting tool, aiming to solve the problems of obvious high-temperature softening and insufficient red hardness of traditional Ni-based or Co-based binder phases in existing Ti(C,N)-based cermets, as well as insufficient wettability, low interfacial bonding strength, difficulty in sintering densification and difficulty in the stable formation of ring core structure in systems with high tungsten content.

[0006] This invention provides a method for preparing a high red hardness Ti(C,N)-based cermet cutting tool material, comprising the following steps: Weigh out 40%–70% Ti(C,N) ceramic powder, 30%–60% pre-alloyed high-density tungsten alloy powder, and 0%–5% grain refiner by mass percentage. The tungsten content in the pre-alloyed high-density tungsten alloy powder is not less than 90%, and the remaining components are at least two of nickel, cobalt, iron, and molybdenum. The grain refiner is at least one of tantalum carbide, niobium carbide, and molybdenum dicarbide. Ti(C,N) ceramic powder, pre-alloyed high-density tungsten alloy powder and grain refiner were added to the ball milling media and wet ball milled to obtain a mixed slurry. The mixed slurry is dried, sieved, and granulated with a forming agent to obtain mixed granules. The mixed granules are pressed into green bodies; The green blank is subjected to vacuum degreasing and vacuum sintering in sequence, which causes selective dissolution of the surface layer of Ti(C,N) particles and precipitation of ring layers around the Ti(C,N) core during the subsequent cooling process. After the sintered blank is cooled to 800°C, it is filled with protective gas and then cooled to room temperature to obtain a high red hardness Ti(C,N)-based cermet cutting tool material with a ring core structure.

[0007] Preferably, the pre-alloyed high-density tungsten alloy powder is prepared by mixing tungsten powder with at least two of nickel powder, cobalt powder, iron powder, and molybdenum powder, followed by high-energy ball milling, pressing, pre-sintering, crushing, and re-ball milling to obtain the pre-alloyed high-density tungsten alloy powder.

[0008] Preferably, the wet ball milling uses anhydrous ethanol or hexane as the milling medium, with a ball-to-material ratio of 4:1 to 6:1, a milling speed of 200 r / min to 300 r / min, and a milling time of 48 h to 72 h.

[0009] Preferably, the drying temperature is 60℃~80℃, the drying time is 6h~24h, the sieve mesh size is 40 mesh~200 mesh, and the forming agent is a polyethylene glycol aqueous solution.

[0010] Preferably, the pressing pressure is 100MPa to 220MPa.

[0011] Preferably, the vacuum degreasing temperature is 500℃~600℃, the holding time is 1h~3h, and the vacuum degree during the degreasing stage is below 10Pa; the vacuum sintering temperature is 1400℃~1550℃, the holding time is 45min~240min, and the vacuum degree during the sintering stage is... .

[0012] Preferably, argon is used as a protective gas in the cooling step.

[0013] This invention also provides a high red hardness Ti(C,N)-based cermet cutting tool material, comprising a Ti(C,N) ceramic phase, a high-density tungsten alloy binder phase, and a grain refiner. The Ti(C,N) ceramic phase comprises 40%–70% by mass, the high-density tungsten alloy binder phase comprises 30%–60% by mass, and the grain refiner comprises 0%–5% by mass. The high-density tungsten alloy binder phase contains at least 90% tungsten, and the remaining components are at least two of nickel, cobalt, iron, and molybdenum. The grain refiner is at least one of tantalum carbide, niobium carbide, and molybdenum dicarbide. The Ti(C,N) ceramic phase comprises a Ti(C,N) core and a ring layer surrounding the Ti(C,N) core. The ring layer is a solid solution layer containing tungsten and the grain refiner element. The high-density tungsten alloy binder phase is distributed between adjacent Ti(C,N) ceramic phases and forms a metallurgical bonding interface with the ring layer.

[0014] The present invention also provides a cutting tool comprising the above-mentioned high red hardness Ti(C,N)-based cermet cutting tool material.

[0015] The beneficial effects of this invention are: This invention uses a high-density tungsten alloy with a tungsten content of not less than 90% as the main binder phase source, and reconstructs the binder phase system at the material matrix level, so that the resulting material can still maintain high strength and good creep resistance under high temperature conditions, thereby improving the red hardness of the cutting tool.

[0016] This invention uses pre-alloyed high-density tungsten alloy powder instead of simply mechanically mixing elemental tungsten powder with other metal powders. This is beneficial for improving component uniformity, reducing local segregation, and improving the uniformity of liquid phase distribution during subsequent sintering, thereby improving the wetting state and densification effect between the Ti(C,N) ceramic phase and the high-density tungsten alloy binder phase.

[0017] This invention constructs a ring-core structure by adding at least one grain refiner selected from tantalum carbide, niobium carbide, and molybdenum dicarbide to the composition, combined with vacuum degreasing, vacuum sintering, and controlled cooling. This causes selective dissolution of the surface layer of Ti(C,N) particles and precipitation of a ring layer around the Ti(C,N) core, thus creating a ring-core structure. This ring-core structure can alleviate the thermal expansion mismatch stress between the ceramic and metallic phases, improve the interfacial metallurgical bonding strength, and generate deflection and hindrance effects during crack propagation, thereby improving transverse fracture strength and resistance to chipping. Attached Figure Description

[0018] Figure 1 This is a schematic flowchart of the method steps of the present invention.

[0019] Figure 2This is a schematic diagram of the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material of the present invention.

[0020] Figure 3 This is a schematic diagram showing the positional relationship between the Ti(C,N) core, the ring layer, and the high-density tungsten alloy binder phase in this invention.

[0021] Figure 4 This is a schematic diagram of the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material obtained in Example 1 of the present invention.

[0022] Figure 5 This is a schematic diagram of the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material obtained in Example 2 of the present invention.

[0023] Figure 6 This is a schematic diagram of the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material obtained in Example 3 of the present invention. Detailed Implementation

[0024] The technical solutions of the present invention will be clearly and completely described below with reference to the accompanying drawings of the embodiments of the present invention. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments.

[0025] like Figure 1 As shown in the embodiment of the present invention, a method for preparing a high red hardness Ti(C,N)-based cermet cutting tool material is provided, comprising the following steps: powder preparation, mixing, granulation, pressing, sintering, and cooling. Wherein: In the powder preparation step, 40%–70% Ti(C,N) ceramic powder, 30%–60% pre-alloyed high-density tungsten alloy powder, and 0–5% grain refiner are weighed by mass percentage. The tungsten content in the pre-alloyed high-density tungsten alloy powder is not less than 90%, and the remaining components are at least two of nickel, cobalt, iron, and molybdenum. The grain refiner is at least one of tantalum carbide, niobium carbide, and molybdenum dicarbide. In the mixing step, the Ti(C,N) ceramic powder, the pre-alloyed high-density tungsten alloy powder and the grain refiner are added to the ball milling media and wet ball milled to obtain a mixed slurry. In the granulation step, the mixed slurry is dried, sieved, and a forming agent is added for granulation to obtain mixed granules; The pressing step involves pressing the mixed particles into a green compact; In the sintering step, the green blank is subjected to vacuum degreasing and vacuum sintering in sequence, so that the surface layer of Ti(C,N) particles is selectively dissolved and a ring layer is precipitated around the Ti(C,N) core during the subsequent cooling process. In the cooling step, the sintered blank is cooled to 800°C, then filled with protective gas and cooled to room temperature to obtain a high red hardness Ti(C,N)-based cermet tool material with a ring core structure.

[0026] The high red hardness Ti(C,N)-based cermet cutting tool material provided by this invention uses a Ti(C,N) ceramic phase as a hard skeleton and a high-density tungsten alloy as the main binder phase source. At least one grain refiner selected from tantalum carbide, niobium carbide, and molybdenum dicarbide is introduced into the system. This system does not simply increase the content of a single component, but rather constructs a stable ring-core structure through pre-alloyed high-density tungsten alloy powder, active metal element-assisted wetting, and selective dissolution and precipitation reactions during vacuum sintering, thereby achieving a balance between high-temperature hardness, interfacial bonding strength, and cutting toughness.

[0027] The preferred mass percentage of the Ti(C,N) ceramic phase is 40%–70%. When the Ti(C,N) ceramic phase content is too low, the material's hardness and wear resistance are insufficient; when the Ti(C,N) ceramic phase content is too high, it is detrimental to sintering densification and weakens the buffering effect of the binder on impact loads. The preferred mass percentage of the high-density tungsten alloy binder phase is 30%–60%. When its content is too low, it is insufficient to form a continuous and effective metal binder network; when its content is too high, it weakens the overall proportion of the hard skeleton. The preferred mass percentage of the grain refiner is 0–5%. Its function is twofold: firstly, to inhibit abnormal grain growth; and secondly, to regulate the surface tension of the liquid phase and the carbon-nitrogen potential during sintering, thereby promoting the formation of the ring-core structure.

[0028] High-density tungsten alloys are preferably multi-component alloys with tungsten as the main component, wherein the tungsten content is not less than 90%, and the remaining components are at least two of nickel, cobalt, iron, and molybdenum. Tungsten mainly provides a high melting point, high temperature modulus, and creep-resistant framework; nickel, cobalt, and iron are mainly used to form a liquid phase that is conducive to wetting and flow within the sintering temperature range; molybdenum is mainly used for solid solution strengthening and to improve the wetting properties of the liquid phase on Ti(C,N). Due to the use of pre-alloyed high-density tungsten alloy powder, it is more beneficial to obtain a uniform liquid phase and a continuous binder phase network during sintering than simple mechanical mixing of elemental tungsten powder with nickel powder and cobalt powder.

[0029] In some embodiments, pre-alloyed high-density tungsten alloy powder can be prepared first. Specifically, tungsten powder can be mixed with at least two of nickel powder, cobalt powder, iron powder, and molybdenum powder in a preset ratio, and then pressed into shape after high-energy ball milling. After pre-sintering at about 1480°C, an alloy block is formed. Subsequently, the alloy block is crushed and ball-milled again to obtain pre-alloyed high-density tungsten alloy powder with a particle size suitable for subsequent composite sintering, preferably with a median particle size of 0.6 μm to 0.9 μm.

[0030] In the mixing step, Ti(C,N) ceramic powder, pre-alloyed high-density tungsten alloy powder, and grain refiner can be added together to a ball mill jar, along with anhydrous ethanol or hexane as the milling medium. Wet ball milling is preferred, with a ball-to-powder ratio of 4:1 to 6:1, a milling speed of 200 r / min to 300 r / min, and a milling time of 48 h to 72 h. For systems with higher ceramic content, the milling time can be appropriately extended to improve mixing uniformity; for ultrafine powder systems, the milling time can be shortened at higher speeds to reduce the risk of agglomeration.

[0031] In the granulation step, the ball-milled slurry can be dried in a vacuum drying oven at 60℃~80℃ for a preferred time of 6h~24h. After drying, it is sieved, preferably with a mesh size of 40 mesh~200 mesh. Subsequently, a polyethylene glycol aqueous solution is added for granulation. The polyethylene glycol is preferably PEG-4000, and the amount added can be 2% of the total mass of the mixed powder to improve the fluidity and formability in the subsequent pressing process.

[0032] In the pressing step, the obtained mixed particles can be pressed into blade green bodies under a pressure of 100MPa to 220MPa. The preferred pressure is around 200MPa; when the ceramic phase content is high, the pressure can be appropriately increased to 220MPa to further improve the green body density. Specifically, a standard blade green body with a diameter of Φ12.7mm × 4.76mm is used.

[0033] In the sintering process, it is preferable to first perform vacuum debinding followed by vacuum sintering. During the vacuum debinding stage, the green compact can be heated to 500℃~600℃ and held for 1h~3h, with a vacuum level below 10Pa, to remove the forming agent. After debinding, the temperature is further increased to 1400℃~1550℃ for vacuum sintering, held for 45min~240min, and the vacuum level is maintained at... For Fe-containing systems, a lower sintering temperature can be used because the liquid phase appears at a lower temperature; for ultrafine-grained systems, a lower temperature and a shorter holding time can be used to suppress grain growth.

[0034] During vacuum sintering, the surface layer of Ti(C,N) particles undergoes selective dissolution, with elements such as Ti, C, and N migrating into the surrounding liquid phase. Meanwhile, W, Mo, and some grain refiner elements in the liquid phase recrystallize again around the Ti(C,N) core during subsequent cooling, forming a ring layer with a compositional gradient. This ring layer, together with the Ti(C,N) core, constitutes a ring-core structure. Because this transition layer is located between the Ti(C,N) core and the high-density tungsten alloy binder phase, it can create a gradient transition at the interface, improving the interfacial metallurgical bonding strength and reducing thermal expansion mismatch stress concentration.

[0035] like Figure 2As shown, the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material of the present invention includes a high-density tungsten alloy binder phase, Ti(C,N) grains, and a ring layer located between the high-density tungsten alloy binder phase and the Ti(C,N) grains. The Ti(C,N) grains constitute the hard supporting framework of the material, the high-density tungsten alloy binder phase is distributed between adjacent Ti(C,N) grains, and the ring layer is formed in the peripheral region of the Ti(C,N) grains, thereby forming a transitional connection structure between the hard phase and the binder phase. This microstructure configuration improves the interfacial bonding between the high-density tungsten alloy binder phase and the Ti(C,N) grains, and is beneficial for improving the material's high-temperature hardness, chipping resistance, and overall cutting performance.

[0036] like Figure 3 As shown, in the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material of the present invention, the Ti(C,N) core is located in the central region of each hard particle, and a ring layer covers the periphery of the Ti(C,N) core. A high-density tungsten alloy binder phase is distributed between adjacent Ti(C,N) cores. Through selective dissolution during vacuum sintering and precipitation during subsequent cooling, a ring layer is formed between the Ti(C,N) core and the high-density tungsten alloy binder phase. This results in a transitional connection structure between the Ti(C,N) core and the high-density tungsten alloy binder phase, rather than a direct abrupt contact. This structure is beneficial for improving the interfacial bonding between the hard phase and the binder phase, alleviating stress concentration caused by thermal expansion mismatch, and improving the strength, toughness, and high-temperature cutting stability of the material.

[0037] In the cooling step, it is preferable to cool the sintered green body to 800°C in the furnace, then purge it with argon gas, and continue cooling to room temperature to lock the ring core structure. For different formulation systems, the ring layer thickness and microstructure uniformity can also be adjusted by staged cooling. For example, the sintering temperature can be reduced from 1200°C at a cooling rate of 20°C / min, and then reduced from 1200°C to 800°C at a cooling rate of 50°C / min. If cooling is too slow, it may lead to coarsening of the precipitated phase or indistinct ring layers.

[0038] Example 1

[0039] This embodiment provides a high-W-content, high-hardness, high-red-hardness Ti(C,N)-based cermet cutting tool material. The following raw materials are weighed by mass percentage: 55% Ti(C,N) powder, 42% high-density tungsten alloy powder, 2% tantalum carbide, and 1% molybdenum dicarbide, with an additional 2% (by mass) of forming agent relative to the total mass of the mixed powder. The high-density tungsten alloy powder is a W-7Ni-3Co system, with a tungsten content of 90%.

[0040] The powder was placed in a cemented carbide ball mill jar, anhydrous ethanol was added as the milling medium, and WC-Co cemented carbide balls were used as the grinding balls. The ball-to-powder ratio was 4:1, the milling speed was 200 r / min, and the milling time was 48 h. After milling, the resulting slurry was dried in a vacuum drying oven at 60℃, passed through an 80-mesh sieve, and then granulated with a 2% PEG-4000 aqueous solution to obtain well-flowing mixed particles. Subsequently, it was pressed into standard cutting blanks with a diameter of Φ12.7 mm × 4.76 mm under a pressure of 200 MPa.

[0041] The green blanks were placed in a vacuum sintering furnace, first heated to 600℃ and held for 1 hour, with the vacuum level below 10 Pa, to complete degreasing; then the temperature was further increased to 1460℃ and held for 60 minutes, with the vacuum level maintained at [missing value]. The furnace was then cooled to 800°C and argon gas was introduced, followed by further cooling to room temperature. The resulting material showed a clear Ti(C,N) core surrounded by a W- and Ta-rich solid solution ring. The high-density tungsten alloy matrix was uniformly distributed with a porosity of less than 0.5%.

[0042] like Figure 4 As shown, the microstructure of the high red hardness Ti(C,N)-based cermet tool material obtained in Example 1 of this invention is relatively uniform and dense. The Ti(C,N) hard phase and the high-density tungsten alloy binder phase are relatively evenly distributed, and no obvious large-area segregation regions are observed. Combined with... Figure 2 and Figure 3 As shown in the organizational relationship, a ring layer is formed around the Ti(C,N) core, thereby forming a transitional connection structure between the Ti(C,N) core and the high-density tungsten alloy binder phase. This is beneficial for improving the interfacial bonding state and enhancing the high-temperature service performance of the material.

[0043] Example 2

[0044] This embodiment provides a high-strength, high-red-hardness Ti(C,N)-based cermet cutting tool material with a matching toughness. The following raw materials are weighed by weight percentage: 65% Ti(C,N) powder, 33% high-density tungsten alloy powder, 1.5% niobium carbide, and 0.5% tantalum carbide, with an additional 2% (by weight of the total powder mixture) forming agent. The high-density tungsten alloy powder is a W-5Ni-3Fe-2Mo system, with a tungsten content of 90%.

[0045] The remaining processes are basically the same as in Example 1, except that the ball milling time is extended to 60 hours, the pressing pressure is increased to 220 MPa, the sintering temperature is adjusted to 1440℃ and held for 90 minutes, and a segmented cooling system is adopted, with a cooling rate of 20℃ / min from 1440℃ to 1200℃ and a cooling rate of 50℃ / min from 1200℃ to 800℃. This implementation method achieves a balance between hardness and toughness under high Ti(C,N) content conditions, making the resulting material more suitable for intermittent cutting conditions.

[0046] like Figure 5 As shown, the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material obtained in Example 2 of this invention is relatively uniform and dense, and the distribution of Ti(C,N) hard phase particles and high-density tungsten alloy binder phase is relatively coordinated. Combined with... Figure 2 and Figure 3 As shown in the microstructure, a ring layer is formed around the Ti(C,N) core, thus creating a transitional connection structure between the Ti(C,N) core and the high-density tungsten alloy binder phase. Because Example 2 employed a high Ti(C,N) ceramic phase content and a segmented cooling process, the resulting microstructure maintains a high proportion of hard phase while still exhibiting good interfacial bonding and microstructure uniformity, which is beneficial for improving the strength and toughness balance of the material.

[0047] Example 3

[0048] This embodiment provides an ultrafine-grained, high-red-hardness Ti(C,N)-based cermet cutting tool material. The following raw materials are weighed by mass percentage: 60% ultrafine Ti(C,N) powder, 38% high-density tungsten alloy powder, and 2% molybdenum dicarbide, with an additional 2% (by mass) of forming agent relative to the total mass of the mixed powder. The high-density tungsten alloy powder is a W-6Ni-4Co system, with a tungsten content of 90%.

[0049] The materials were mixed using a planetary ball mill at a speed of 300 rpm for 36 hours. After drying and granulation, the mixture was pressed into green bodies under a pressure of 200 MPa. After degreasing, the sintering temperature was lowered to 1420℃ and held for 45 minutes to inhibit grain growth. During the holding stage, 0.5 MPa of high-purity argon gas was introduced to promote pore closure. The resulting material has a grain size of less than 1 μm, with thin and dense ring layers, making it suitable for high-precision machining applications.

[0050] like Figure 6 As shown, the microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material obtained in Example 3 of this invention is relatively fine and uniform, with significantly refined Ti(C,N) hard phase particles dispersed within a high-density tungsten alloy binder phase. Combined with... Figure 2 and Figure 3As shown in the microstructure, a ring layer is formed around the Ti(C,N) core, thus creating a transitional bonding structure between the Ti(C,N) core and the high-density tungsten alloy binder phase. Because Example 3 uses ultrafine Ti(C,N) powder and employs a lower sintering temperature and shorter holding time, it is possible to maintain a good interfacial bonding state while suppressing abnormal grain growth, thereby improving the material's hardness and machinability.

[0051] Comparative Example 1 This comparative example uses a traditional Ni binder system, weighing out 60% Ti(C,N) powder, 37% Ni powder, 2% tantalum carbide, and 1% molybdenum dicarbide by mass percentage, and vacuum sintering at 1400℃ for 60 min. The resulting material has no obvious ring-core structure, and the matrix softens significantly at high temperatures.

[0052] Comparative Example 2 This comparative example uses a mechanical mixing route for elemental tungsten powder. The powder composition is as follows: 60% Ti(C,N) powder, 30% elemental W powder, 7% Ni powder, and 3% Co powder by mass percentage. The overall composition is similar to that of Example 1, but pre-alloyed high-density tungsten alloy powder was not used. The remaining processes are the same as in Example 1. The resulting material shows incomplete dissolution of W particles, numerous unmelted nuclei and pores, weak interfacial bonding, and significant component segregation.

[0053] Comparative Example 3 The formulation of this comparative example is basically the same as that of Example 1, but tantalum carbide and molybdenum dicarbide are not added, and the furnace is cooled at an extremely slow rate of less than 5°C / min after sintering. The resulting material has an indistinct ring-core structure, coarser grains, straight interfaces, and cracks are more likely to propagate along the interfaces.

[0054] Performance testing methods The above-described embodiments and comparative samples were subjected to wire cutting, grinding, and polishing before performance testing. Density was measured using the Archimedes water displacement method; hardness was measured using a Rockwell hardness tester, with high-temperature hardness tested at 1000℃; transverse fracture strength was tested using the three-point bending method; fracture toughness was tested using the indentation method; high-temperature oxidation resistance was evaluated by the weight gain per unit area after static air oxidation at 1000℃ for 2 hours; cutting life was evaluated by dry turning of Inconel 718 high-temperature alloy with cutting parameters of 120 m / min, feed rate of 0.15 mm / r, and depth of cut of 1.0 mm. The failure criterion was a flank wear VB reaching 0.3 mm.

[0055] Table 1 Comparison of performance parameters between each embodiment and the comparative example.

[0056] As shown in Table 1, the embodiments of the present invention are superior to the comparative examples in terms of high-temperature hardness at 1000℃, oxidation weight gain, and cutting life. Specifically, the high-temperature hardness at 1000℃ in Example 1 is 86.5 HRA, and the oxidation weight gain at 1000℃ is 0.08 mg / cm³. 2 The cutting life was 42 min, while the corresponding data for Comparative Example 1 were 78.2 HRA and 0.45 mg / cm², respectively. 2 The results were 18 min and 18 min respectively; Comparative Examples 2 and 3 also showed problems of weak interfacial bonding or unclear ring core structure. This indicates that the present invention, through the design of a pre-alloyed high-density tungsten alloy system, control of grain refiner, and construction of a vacuum sintered ring core structure, can effectively improve the overall service performance of Ti(C,N)-based cermet cutting tool materials.

[0057] As shown in Table 1, the transverse fracture strength of Example 2 is 2680 MPa, and the fracture toughness is 12.8 MPa·m. 1 / 2 All of these levels are superior to the corresponding levels in Examples 1 and 3. Figure 5 The microstructure shown indicates that Example 2, even with a high Ti(C,N) ceramic phase content, still maintains a relatively uniform microstructure distribution and good interfacial bonding. This demonstrates that the present invention, through the design of a high-density tungsten alloy system, the regulation of grain refiners, and the control of segmented cooling processes, helps to achieve a better match between material strength and toughness.

[0058] As shown in Table 1, the room temperature hardness of Example 3 is 93.8 HRA and the high temperature hardness at 1000℃ is 87.2 HRA, which are all at a relatively high level among the three examples. Figure 6 The microstructure shown indicates that the material obtained in Example 3 has a finer microstructure, which corresponds to its process route of using ultrafine Ti(C,N) powder, lower sintering temperature and shorter holding time. This shows that the present invention, through fine-grained microstructure control, is beneficial to improving the material hardness and edge retention performance under high-temperature cutting conditions.

[0059] Figures 2-6 In detail: Figure 2 The microstructure of the high red hardness Ti(C,N)-based cermet cutting tool material of the present invention is shown, wherein point A is a ring phase interface (i.e., a ring layer), point B is Ti(C,N) grains, and point C is a high-density alloy phase (i.e., a high-density tungsten alloy binder phase). Figure 2It can be seen that the microstructure of this material is not a single homogeneous structure, but rather consists of Ti(C,N) grains, a high-density tungsten alloy binder phase, and ring layers. In the figure, the Ti(C,N) grains are distributed within the material and form a hard supporting framework; the transition region surrounding the Ti(C,N) grains corresponds to the ring layers; and the continuous distribution region between adjacent Ti(C,N) grains corresponds to the high-density tungsten alloy binder phase. Therefore, the material obtained by this invention is not a simple two-phase composite structure in the traditional sense where the hard phase and binder phase are in direct contact, but rather an intermediate transition layer is constructed between the hard phase and the binder phase.

[0060] Further analysis Figure 2 As can be seen from the microstructure, the ring layer is located between the Ti(C,N) grains and the high-density tungsten alloy binder phase. Its function is to create a gradual transition in the microstructure between the Ti(C,N) grains and the high-density tungsten alloy binder phase from the inside out, rather than forming an abrupt contact in terms of interface composition and properties. This microstructure configuration is beneficial for improving the coating and bonding state of the high-density tungsten alloy binder phase with the Ti(C,N) grains, and for reducing interfacial stress concentration caused by differences in thermal expansion coefficients within the material after sintering. In other words, Figure 2 This does not reflect a simple phase coexistence state, but rather a three-layered composite structure of hard phase, ring layer and binder phase established by the present invention through component design and sintering control.

[0061] From the perspective of material properties, Figure 2 The microstructure shown is advantageous for simultaneously achieving high hardness and high-temperature stability. Specifically, the Ti(C,N) grains ensure high hardness and wear resistance, while the high-density tungsten alloy binder phase guarantees good strength and creep resistance under high-temperature conditions. The annular layer between the two provides a microstructure transition at the interface. Therefore, Figure 2 The microstructure shown is an important microstructure basis for achieving high red hardness, high interfacial bonding strength and good overall cutting performance in this invention.

[0062] Figure 3 The positional relationship between the Ti(C,N) core, the ring layer, and the high-density tungsten alloy binder phase in this invention is shown. Figure 3 (a) Schematic representation of the microstructure evolution under the combined action of Ti(C,N) particles, high-density tungsten alloy phase, and refining agent during sintering. Figure 3 (b) Schematic representation of the microstructure of the ring layers formed after sintering and controlled cooling. From Figure 3 (a) It can be seen that the Ti(C,N) particles come into contact with and interact with the high-density tungsten alloy phase and the refining agent, and an interfacial transition region gradually forms around the particles; from Figure 3(b) It can be seen that in the final structure, the Ti(C,N) core is located in the central region, the ring layer covers the outer periphery of the Ti(C,N) core, and the high specific gravity tungsten alloy binder phase is distributed on the outer side of the ring layer and between adjacent particles.

[0063] Figure 3 Intuitively, the microstructure of this invention is not one where the Ti(C,N) core is directly exposed within the high-density tungsten alloy binder phase. Instead, the Ti(C,N) core first forms ring layers outwards, and then these ring layers connect with the high-density tungsten alloy binder phase, thus forming a hierarchical microstructure of Ti(C,N) core-ring layers-high-density tungsten alloy binder phase. This hierarchical relationship means that within the material, the innermost layer is the Ti(C,N) core responsible for high hardness, the middle layer is the ring layers responsible for transitional bonding, and the outermost layer is the high-density tungsten alloy binder phase responsible for bonding and high-temperature support. The resulting microstructure is not a simple mixture, but a composite structure with clear spatial and interfacial division of labor.

[0064] Further analysis of the process mechanism, Figure 3 The illustrated microstructure corresponds to the selective dissolution and subsequent cooling precipitation processes that occur during vacuum sintering in this invention. Specifically, some elements in the surface layer of the Ti(C,N) particles migrate to the surrounding liquid phase during sintering, and subsequently, W elements and refining agent-related elements in the liquid phase reprecipitate around the Ti(C,N) core during cooling, thereby forming a ring layer. Therefore, Figure 3 This not only demonstrates the static organizational positional relationship, but also reflects the technical concept of the present invention to achieve in-situ formation of ring layers through process control.

[0065] From the perspective of material service Figure 3 The microstructure of the Ti(C,N) core-ring layer-high-density tungsten alloy binder phase shown is conducive to transforming the interfacial bonding from abrupt contact to gradient transition, thereby improving the interfacial bonding state between the Ti(C,N) core and the high-density tungsten alloy binder phase, alleviating local stress concentration caused by thermal expansion mismatch, and inhibiting rapid crack propagation along the interface. Therefore, Figure 3 This is an important schematic diagram illustrating the rationality of the structure of the present invention and its toughening and stabilizing mechanism.

[0066] Figure 4 The microstructure of the high red hardness Ti(C,N)-based cermet tool material obtained in Example 1 of the present invention is shown. Figure 4It can be seen that the material obtained in Example 1 has a relatively uniform overall microstructure distribution. No obvious large-area segregation regions or significant large-sized abnormal pores were observed inside the material, indicating that this example can obtain a relatively dense and uniform microstructure under the adopted formulation and sintering regime. The distribution of different phase regions in the figure is relatively clear, indicating that the Ti(C,N) hard phase and the high-density tungsten alloy binder phase have a good composite state inside the material.

[0067] Based on the specific process conditions of Example 1, Example 1 used W-7Ni-3Co system high specific gravity tungsten alloy powder, added tantalum carbide and molybdenum dicarbide, and adopted vacuum sintering at 1460℃ and subsequent argon cooling process. Figure 4 The uniformity of the microstructure shown indicates that this process combination enables the high-density tungsten alloy binder phase to be more fully distributed between the Ti(C,N) hard phases, and forms a relatively stable transition region around the Ti(C,N) particles. In other words, Figure 4 The microstructure reflected is not the result of simple sintering densification, but rather the result of the combined effects of a pre-alloyed high-density tungsten alloy system, refining agents, and controlled sintering cooling.

[0068] Further integration Figure 2 and Figure 3 Organizational relationship analysis, Figure 4 A corresponding annular region should be formed between the uniformly distributed Ti(C,N) hard phase and the high-density tungsten alloy binder phase, thereby constructing a stable transitional connection structure between the Ti(C,N) core and the high-density tungsten alloy binder phase. This microstructure means that Example 1 can better balance the high proportion of hard phase and the relatively stable distribution of binder phase, which is beneficial for maintaining good interfacial bonding strength and high-temperature service stability while ensuring high hardness.

[0069] From the performance correlation, Example 1 exhibits higher high-temperature hardness at 1000℃, lower oxidation weight gain, and longer cutting life as shown in Table 1. Figure 4 The uniform and relatively dense microstructure shown is a crucial structural support for these performance characteristics. Therefore, Figure 4 This demonstrates the basis for the achievement of Example 1 in terms of high-temperature red hardness and tissue stability.

[0070] Figure 5 The microstructure of the high red hardness Ti(C,N)-based cermet tool material obtained in Example 2 of the present invention is shown. Figure 5 It can be seen that the phase distribution within the material obtained in Example 2 is still relatively uniform, the overall microstructure is relatively dense, and the Ti(C,N) hard phase particles and the high-density tungsten alloy binder phase maintain a relatively coordinated distribution. Compared with Example 1, Figure 5The higher proportion of hard phase particles in the structure shown indicates that Example 2, after increasing the proportion of Ti(C,N) ceramic phase, can still maintain good structural uniformity and interphase composite state.

[0071] As can be seen from the formulation of Example 2, this example uses a high proportion of Ti(C,N) powder and W-5Ni-3Fe-2Mo system high specific gravity tungsten alloy powder, and adds niobium carbide and tantalum carbide, while adopting a segmented cooling system. Figure 5 The results show that even under conditions with a high proportion of hard phase, by designing a high-density tungsten alloy system containing Fe and Mo and controlling the segmented cooling process, the high-density tungsten alloy binder phase can still be relatively evenly distributed inside the material, and significant microstructure imbalance caused by the increase in the proportion of hard phase can be avoided.

[0072] Further analysis of interface organization, combined with Figure 2 and Figure 3 It can be seen that, Figure 5 In the microstructure of Example 2, a ring layer still forms around the Ti(C,N) core, thus establishing a transitional connection structure between the Ti(C,N) core and the high-density tungsten alloy binder phase. Because Example 2 uses a higher proportion of Ti(C,N) ceramic phase, theoretically, it is more prone to problems such as increased contact between hard phases, insufficient binder phase, and interface instability. Figure 5 The presence of good microstructure uniformity indicates that the high-density tungsten alloy system and process route used in this invention have a significant effect on maintaining interfacial stability under conditions of high hard phase ratio.

[0073] According to the performance analysis in Table 1, Example 2 exhibits superior transverse fracture strength and fracture toughness among the three examples, reaching 2680 MPa and 12.8 MPa·m, respectively. 1 / 2 . Figure 5 The relatively uniform and harmonious microstructure shown indicates that Example 2, while increasing the proportion of the hard phase, still maintains good interphase bonding and microstructure continuity, thus being more conducive to achieving a better balance between strength and toughness. In other words, Figure 5 This reflects the microstructure of Example 2, which maintains a good strength-toughness match even under conditions of higher hard phase content.

[0074] Figure 6 The microstructure of the high red hardness Ti(C,N)-based cermet tool material obtained in Example 3 of the present invention is shown. Figure 6As can be seen, the microstructure of the material obtained in Example 3 is relatively fine and uniformly distributed, and no obvious large-sized abnormal grain aggregation regions are observed inside the material, indicating that this example has a good effect in controlling the microstructure refinement. The hard phase particles in the figure are relatively smaller, reflecting that Example 3 has obvious fine-grained microstructure characteristics.

[0075] Based on the analysis of the process conditions in Example 3, Example 3 used ultrafine Ti(C,N) powder, a lower sintering temperature, and a shorter holding time, and introduced high-purity argon gas during the holding stage to promote pore closure. Figure 6 The finer microstructure shown illustrates that by synergistically controlling the raw material grain size and sintering process, abnormal grain growth can be effectively suppressed, maintaining a finer microstructure in the material. In other words, Figure 6 This does not reflect the random grain refinement phenomenon under ordinary sintering conditions, but rather the result of fine grain structure control achieved by the present invention through the combined use of raw material system and heat treatment regime.

[0076] Further integration Figure 2 and Figure 3 Analysis of the microstructure in Example 3 showed that while maintaining a refined microstructure, a ring layer was still formed around the Ti(C,N) core, ensuring a smooth transition between the Ti(C,N) core and the high-density tungsten alloy binder phase. This is crucial because a lack of stable interface transitions in a fine-grained structure can easily lead to instability at the interface around the fine particles or local microstructural imbalance. Figure 6 The relatively uniform and refined microstructure shown indicates that the present invention does not simply pursue grain refinement, but maintains a good interphase bonding state on the basis of refinement.

[0077] Based on the performance correlation, Example 3 exhibits the highest room temperature hardness and a relatively high high temperature hardness of 1000℃ in Table 1. Figure 6 The refined and uniform microstructure shown demonstrates that Example 3, through a process route using ultrafine Ti(C,N) powder, a lower sintering temperature, and a shorter holding time, achieved a microstructure beneficial to improving material hardness and edge retention. Therefore, Figure 6 This invention primarily demonstrates its advantages in terms of high hardness and applicability to fine machining.

[0078] In conclusion, Figure 2 and Figure 3 This invention mainly describes the microstructure and hierarchical relationship of the high red hardness Ti(C,N)-based cermet cutting tool material, namely, a transitional connection structure formed by a Ti(C,N) core, ring layers, and a high specific gravity tungsten alloy binder phase. Figures 4 to 6The actual microstructure characteristics of the materials obtained under different embodiment conditions are described respectively. Among them, Example 1 shows a more uniform and dense microstructure, Example 2 shows a good strength-toughness balance even under a higher proportion of hard phase, and Example 3 shows a significant advantage in fine-grained microstructure. This shows that the present invention can not only form a stable Ti(C,N) core-ring layer-high specific gravity tungsten alloy binder phase microstructure, but also optimize the performance of the material in different aspects such as high hardness, high strength-toughness balance, and applicability to fine-grained finishing by adjusting the formulation composition and sintering process.

[0079] Compared with the prior art, the present invention has at least the following beneficial effects.

[0080] This invention uses pre-alloyed high-density tungsten alloy powder as the main binder phase source in the powder preparation step, ensuring that the tungsten content in the high-density tungsten alloy is not less than 90%, with the remaining components being at least two of nickel, cobalt, iron, and molybdenum. This technique does not rely on traditional Ni or Co as the main binder phase; instead, it reconstructs the binder phase system using a high-density tungsten alloy, giving the binder phase itself a higher melting point, high-temperature strength, and creep resistance. This improves the thermal stability and red hardness of the cutting tool under high-temperature cutting conditions from the material matrix level. Performance data shows that the sample in the example exhibits superior hardness at 1000℃, oxidation weight gain, and cutting life compared to samples with a traditional Ni binder phase, demonstrating the clear technical effectiveness of this binder phase system replacement.

[0081] This invention employs pre-alloyed high-density tungsten alloy powder, instead of directly mechanically mixing and sintering elemental W powder with Ni and Co powders. This technique helps improve the pre-homogenization of the tungsten components and other metal components, reducing problems such as unmelted nuclei, segregation, and interfacial discontinuities caused by insufficient local component dispersion during subsequent sintering. This improves the wetting state of the high-density tungsten alloy binder relative to the Ti(C,N) ceramic phase and the densification effect during sintering. Comparative Example 2 has shown that without the pre-alloying route, problems such as incomplete dissolution of W particles, numerous pores, and weak interfacial bonding occur. Therefore, including pre-alloyed high-density tungsten alloy powder in the independent claims has strong technical support.

[0082] This invention incorporates at least one grain refiner selected from tantalum carbide, niobium carbide, and molybdenum dicarbide into the composition. During the sintering process, vacuum degreasing, vacuum sintering, and subsequent controlled cooling selectively dissolve the surface layer of Ti(C,N) particles, precipitating a ring layer around the Ti(C,N) core to obtain a ring-core structure. This technique does not simply pursue dense sintering, but rather utilizes the combination of the composition system and sintering regime to induce the formation of a solid solution ring layer around the Ti(C,N) core, creating a gradient transition interface between the ceramic phase and the binder phase. This gradient transition interface can mitigate stress concentration caused by thermal expansion mismatch, enhance the metallurgical bond between the two phases, and deflect and impede crack propagation, thus improving the transverse fracture strength and spalling resistance of the material. Comparative Example 3 shows that without the corresponding additives and cooling regime, the ring-core structure is not obvious, the grains are coarse, and cracks easily propagate along the interface, while the examples exhibit higher strength, toughness, and cutting life.

[0083] This invention, through the synergistic combination of the aforementioned component system, pre-alloying route, and sintering control, enables the resulting high-red-hardness Ti(C,N)-based cermet cutting tool material to simultaneously possess high room-temperature hardness, high 1000℃ high-temperature hardness, low oxidation weight gain, and long cutting life. Therefore, it is well-suited for high-speed cutting, dry cutting, and cutting of difficult-to-machine materials. This beneficial effect cannot be directly obtained by changing a single component, but rather is the comprehensive result of the combined effects of a high-density tungsten alloy binder phase, grain refiner, and ring-core structure.

[0084] Finally, it should be noted that the above description is only a preferred embodiment of the present invention and is not intended to limit the present invention. Although the present invention has been described in detail with reference to the foregoing embodiments, those skilled in the art can still modify the technical solutions described in the foregoing embodiments or make equivalent substitutions for some of the technical features. Any modifications, equivalent substitutions, improvements, etc., made within the spirit and principles of the present invention should be included within the protection scope of the present invention.

Claims

1. A method for preparing a high red hardness Ti(C,N)-based cermet cutting tool material, characterized in that, Includes the following steps: In the powder preparation step, 40%–70% Ti(C,N) ceramic powder, 30%–60% pre-alloyed high-density tungsten alloy powder, and 0–5% grain refiner are weighed according to the mass percentage. The tungsten content in the pre-alloyed high-density tungsten alloy powder is not less than 90%. In the mixing step, the Ti(C,N) ceramic powder, the pre-alloyed high-density tungsten alloy powder and the grain refiner are added to the ball milling media and wet ball milled to obtain a mixed slurry. In the granulation step, the mixed slurry is dried, sieved, and a forming agent is added for granulation to obtain mixed granules; The pressing step involves pressing the mixed particles into a green compact; In the sintering step, the green blank is subjected to vacuum degreasing and vacuum sintering in sequence, so that the surface layer of Ti(C,N) particles is selectively dissolved and a ring layer is precipitated around the Ti(C,N) core during the subsequent cooling process. In the cooling step, the sintered blank is cooled to 800°C, then filled with protective gas and cooled to room temperature to obtain a high red hardness Ti(C,N)-based cermet tool material with a ring core structure.

2. The method for preparing the high red hardness Ti(C,N)-based cermet cutting tool material according to claim 1, characterized in that, The powder preparation step further includes: the remaining components are at least two of nickel, cobalt, iron, and molybdenum, and the grain refiner is at least one of tantalum carbide, niobium carbide, and molybdenum dicarbide; The pre-alloyed high-density tungsten alloy powder is prepared by mixing tungsten powder with at least two of nickel powder, cobalt powder, iron powder, and molybdenum powder, followed by high-energy ball milling, pressing, pre-sintering, crushing, and re-ball milling to obtain the pre-alloyed high-density tungsten alloy powder. The pre-sintering temperature is 1450℃~1500℃; The median particle size of the pre-alloyed high-density tungsten alloy powder is 0.6 μm to 0.9 μm.

3. The method for preparing high red hardness Ti(C,N)-based cermet cutting tool material according to claim 2, characterized in that, In the mixing step, the ball milling medium is anhydrous ethanol or hexane, the ball-to-material ratio is 4:1 to 6:1, the ball milling speed is 200 r / min to 300 r / min, and the ball milling time is 48 h to 72 h. In the granulation step, the drying temperature is 60℃~80℃, the drying time is 6h~24h, the sieve mesh size is 40 mesh~200 mesh, and the forming agent is a polyethylene glycol aqueous solution.

4. The method for preparing the high red hardness Ti(C,N)-based cermet cutting tool material according to claim 3, characterized in that, The forming agent is a polyethylene glycol 4000 aqueous solution, and the amount added is 2% of the total mass of the mixed powder.

5. The method for preparing the high red hardness Ti(C,N)-based cermet cutting tool material according to claim 4, characterized in that, The pressing pressure in the pressing step is 100MPa~220MPa; The vacuum degreasing temperature in the sintering step is 500℃~600℃, the holding time is 1h~3h, and the vacuum degree in the degreasing stage is less than 10Pa. The vacuum sintering temperature in the sintering step is 1400℃~1550℃, the holding time is 45min~240min, and the vacuum degree during the sintering stage is... ; Argon is used as the protective gas in the cooling step.

6. The method for preparing the high red hardness Ti(C,N)-based cermet cutting tool material according to claim 5, characterized in that, The cooling step includes: a first cooling stage from sintering temperature to 1200°C, and a second cooling stage from 1200°C to 800°C, wherein the cooling rate of the first cooling stage is 20°C / min, and the cooling rate of the second cooling stage is 50°C / min. In the grain refiner, the amount of tantalum carbide added is 0.5% to 2%, the amount of niobium carbide added is 0.5% to 1.5%, and the amount of molybdenum dicarbide added is 0.5% to 2%.

7. The method for preparing the high red hardness Ti(C,N)-based cermet cutting tool material according to claim 6, characterized in that, In the sintering step, the Ti, C and N elements in the surface layer of Ti(C,N) particles selectively dissolve into the surrounding liquid phase, and together with the W element and grain refiner elements in the liquid phase, precipitate to form the ring layer during the cooling process.

8. A high red hardness Ti(C,N)-based cermet cutting tool material, characterized in that, The mixture comprises a Ti(C,N) ceramic phase, a high-density tungsten alloy binder phase, and a grain refiner. The Ti(C,N) ceramic phase comprises 40%–70% by mass, the high-density tungsten alloy binder phase comprises 30%–60% by mass, and the grain refiner comprises 0%–5% by mass. The high-density tungsten alloy binder phase contains at least 90% tungsten, and the remaining components are at least two of nickel, cobalt, iron, and molybdenum. The grain refiner is at least one of tantalum carbide, niobium carbide, and molybdenum dicarbide. The Ti(C,N) ceramic phase comprises a Ti(C,N) core and a ring layer surrounding the Ti(C,N) core. The ring layer is a solid solution layer containing tungsten and the grain refiner element. The high-density tungsten alloy binder phase is distributed between adjacent Ti(C,N) ceramic phases and forms a metallurgical bonding interface with the ring layer.

9. The high red hardness Ti(C,N)-based cermet cutting tool material according to claim 8, characterized in that, The ring layer surrounds the Ti(C,N) core and forms a gradient transition structure; The thickness of the annular layer is 0.2 μm to 1.5 μm; The porosity of the high red hardness Ti(C,N)-based cermet tool material is no higher than 0.5%; The high-density tungsten alloy binder phase is continuously distributed between adjacent Ti(C,N) ceramic phases.

10. A cutting tool, characterized in that, Including the high red hardness Ti(C,N)-based cermet cutting tool material as described in claim 8 or 9.

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