Binding-phase-free hard alloy and preparation method thereof

Abstract

The invention relates to the technical field of powder metallurgy materials, and discloses a binding-phase-free hard alloy and a preparation method thereof.The binding-phase-free hard alloy comprises, by mass, 75-80 parts of WC; 8 to 12 parts of TiC; 5 to 8 parts of TaC; 3 to 5 parts of ZrC; 1.5 to 2.5 parts of HfC; 0.8 to 1.2 parts of rare earth oxide; the preparation method comprises the five key steps of core-shell structure composite powder preparation, high-energy beam melting gradient accumulation, microwave-magnetic field composite field activation, multi-stage pressure plasma sintering and staged variable-rate gradient cooling. And component gradient design is cooperated with a multi-field coupling process. Through collaborative design of a core-shell structure hard phase and a gradient composite phase, and in combination with a multi-field coupling activated sintering and gradient cooling process, the problem of obdurability balance of the non-binding-phase hard alloy is solved, the microstructure stability is optimized, and the comprehensive performance and the process controllability are improved.

Description

Technical Field

[0001] This invention relates to the field of powder metallurgy materials technology, and in particular to a binderless cemented carbide and its preparation method. Background Technology

[0002] Traditional cemented carbides rely on metallic binders such as cobalt to balance the brittleness of the hard phase. However, the presence of the binder leads to decreased high-temperature strength and limited corrosion resistance. With increasing demands on material performance under extreme conditions, binderless cemented carbides have become an important development direction. In existing technologies, binderless materials prepared by direct sintering of a single superhard phase (such as WC) suffer from difficulties in densification, significant grain boundary brittleness, and are prone to microcracks and other defects during sintering. Some studies have attempted to introduce transition metal carbides as composite phases, but due to unreasonable composition design, the interfacial bonding strength between the hard phases is insufficient, leading to interfacial delamination failure under complex stress.

[0003] In terms of manufacturing processes, conventional sintering methods struggle to achieve sufficient densification of binderless systems, often requiring ultra-high sintering temperatures or ultra-high pressure conditions, leading to severe equipment wear and limited product dimensions. Existing gradient sintering processes mostly employ simple linear compositional changes, failing to effectively coordinate thermal stress distribution, making large-sized products prone to warping and deformation. While novel processes such as microwave sintering can improve the densification process, they impose specific requirements on the material's dielectric properties, have narrow process windows, and struggle to stably control the evolution of the microstructure.

[0004] Therefore, this invention proposes a binderless cemented carbide and its preparation method to overcome the shortcomings of the prior art. Summary of the Invention

[0005] The purpose of this invention is to provide a binderless cemented carbide and its preparation method, which solves the problems of densification difficulties, low interfacial bonding strength, and uneven thermal stress distribution in binderless cemented carbide.

[0006] To achieve the above objectives, the present invention provides the following technical solution: a binderless cemented carbide comprising the following components in parts by mass: WC: 75-80 servings; TiC: 8-12 parts; TaC: 5-8 servings; ZrC: 3-5 parts; HfC: 1.5-2.5 parts; Rare earth oxides: 0.8-1.2 parts.

[0007] Preferably, the WC and TiC form a core-shell structure, the WC core has a particle size of 180-220 nm, the TiC shell has a thickness of 18-22 nm, the surface of the core-shell structure is coated with a rare earth oxide layer with a thickness of 10-15 nm, and a 2-3 nm transition interface layer is formed between the TiC shell and the rare earth oxide layer.

[0008] A core-shell structure is formed by coating WC particles with a TiC shell. Utilizing the lattice matching between TiC and WC (both are hexagonal with a lattice constant difference of <2%), strong bonding is achieved through interfacial diffusion during sintering. The rare earth oxide coating reduces interparticle friction through surface adsorption, while rare earth elements segregate towards grain boundaries at high temperatures, forming nanoscale oxide particles that pin the grain boundaries and inhibit abnormal grain growth. The transition interface layer originates from the interdiffusion of TiC and rare earth oxides during deposition. This submicron-sized transition layer alleviates internal stress caused by differences in thermal expansion coefficients.

[0009] Preferably, the rare earth oxides include Y2O3, La2O3 and Sc2O3, and the molar ratio of the mixture of Y2O3, La2O3 and Sc2O3 is 6.5-7.5:2.5-3.5:0.4-0.6. The rare earth oxides exist on the surface of the core-shell structure in the form of a continuous coating layer with a coating layer thickness of 10-15 nm. The particle size distribution of Y2O3, La2O3 and Sc2O3 in the coating layer is D90≤200 nm.

[0010] Preferably, the TaC, ZrC and HfC form a continuous gradient distribution along the thickness direction, wherein the TaC content decreases from the material surface to the interior according to a quadratic function law, with the surface content decreasing from 100% to 0%, the ZrC content increasing from 0% to 80%, and the HfC content remaining constant at 20%, and the total thickness of the gradient layer is 15-25% of the material thickness.

[0011] A continuous gradient distribution of TaC-ZrC-HfC is constructed along the thickness direction, leveraging the synergistic effect of TaC's high-temperature stability (melting point of 3880℃) and ZrC's high hardness (Vickers hardness of 25.5 GPa). The quadratic decreasing TaC distribution curve forms a high-melting-point protective layer on the material surface, while the increasing ZrC content internally enhances overall toughness. A constant 20% HfC content acts as a structural stabilizer, and its high thermal conductivity of 20 W / m·K optimizes thermal stress distribution. The design of a gradient layer thickness of 15-25% balances functional requirements with process feasibility.

[0012] This invention also provides a method for preparing a binderless cemented carbide, comprising the following steps: S1. Core-shell structure preparation: A continuous titanium carbide shell layer is formed on the surface of tungsten carbide powder by vapor deposition to obtain a core-shell structured composite powder. S2. Gradient layer construction: The core-shell structured composite powder and gradient phase powder are stacked layer by layer using high-energy beam melting technology to form a composition gradient distribution that changes continuously along the thickness direction; S3. Synergistic activation treatment: Under the synergistic effect of the combined field of microwave and alternating magnetic field, the gradient structure blank is activated across scales. S4. Multi-field coupling sintering: Plasma activation sintering with multi-stage pressure loading is applied to the activated green body to achieve material densification and microstructure stabilization. S5. Gradient cooling treatment: The sintered body is subjected to stress regulation by a staged variable rate cooling method to release the residual stress gradient.

[0013] High-energy beam melting: The laser wavelength of 1060-1080nm matches the light absorption characteristics of the WC-TiC system. The absorption rate in this band is >85%. The interlayer rotation angle of 66-70° makes the direction of the molten pool follow the golden ratio, eliminating anisotropy.

[0014] Microwave-magnetic field activation: 2.4-2.5GHz microwaves excite the dielectric loss of the material, and 48-52kHz alternating magnetic fields induce eddy currents to generate Joule heating. The two work together to achieve cross-scale activation of the preform from molecular vibration to lattice vibration.

[0015] Multi-stage pressure sintering: The staged loading rate of 4.5-5.5 → 0.8-1.2 → 0.4-0.6 MPa / min matches the three stages of the material densification process: particle rearrangement → plastic flow → diffusion creep, avoiding grain boundary cracks caused by premature application of high pressure.

[0016] Gradient cooling control: During the rapid cooling phase, the temperature is 45-55℃ / min, which skips the carbide precipitation sensitive temperature range of 1200-1000℃. During the slow cooling phase, the temperature is 8-12℃ / min, which promotes the gradual release of residual stress through dislocation slip.

[0017] Preferably, in step S1: The vapor deposition method is pulsed chemical vapor deposition, with the deposition temperature controlled at 830-870℃, the volume ratio of hydrogen to methane in the carrier gas being 8.5-9.5:1, and the pulse period being 3-5 times. After deposition, the titanium carbide shell thickness in the core-shell structured composite powder is 18-22nm, and the surface roughness Ra of the shell is ≤15nm.

[0018] Preferably, in step S2: The high-energy beam melting technology uses a laser beam with a wavelength of 1060-1080nm, a laser power of 180-220W, a scanning speed of 800-1200mm / s, a single-layer stacking thickness of 45-55μm, an interlayer rotation angle of 66°-70°, and an oxygen content of ≤50ppm in the protective atmosphere.

[0019] Preferably, in step S3: The microwave frequency is 2.4-2.5 GHz, the alternating magnetic field frequency is 48-52 kHz, the magnetic field strength is 0.45-0.55 T, the processing time is 28-32 min, and the temperature gradient of the billet is controlled within 15-25 ℃ / mm during the processing.

[0020] Preferably, in step S4: The multi-stage pressure loading includes three stages: The first stage involves loading at a rate of 4.5-5.5 MPa / min within a pressure range of 0-30 MPa. The second stage involves loading at a rate of 0.8-1.2 MPa / min within the range of 30-45 MPa; The third stage involves loading at a rate of 0.4-0.6 MPa / min within the range of 45-50 MPa; The sintering temperature is maintained at 1450-1550℃, and the holding time is 18-22 minutes.

[0021] Preferably, in step S5: The gradient cooling includes a first stage of cooling from 1600℃ to 1000℃ at a rate of 45-55℃ / min, a second stage of cooling from 1000℃ to 300℃ at a rate of 8-12℃ / min, and a third stage of natural cooling to room temperature, maintaining a vacuum degree ≤5×10⁻⁵ throughout the cooling process. -3 Pa.

[0022] In summary, the present invention has at least one of the following beneficial technical effects: 1. This invention achieves a balance between strength and toughness in materials by synergistically combining an innovatively designed core-shell structure hard phase with a gradient-distributed TaC / ZrC composite phase, while completely removing the metallic binder phase. The core-shell structure effectively alleviates interfacial stress between the hard phases, and the gradient composite phase forms a continuously transitioning reinforcing network, overcoming the technical bottlenecks of traditional binderless cemented carbides, which are characterized by high brittleness and poor machinability.

[0023] 2. This invention employs a rare-earth oxide synergistic doping strategy to form a nanocrystalline grain boundary phase during sintering. This design not only suppresses abnormal grain growth but also enhances grain boundary bonding strength through interfacial electronic structure modulation, solving the industry-wide problem of grain boundary embrittlement in binderless materials.

[0024] 3. This invention achieves synergistic control of powder particle orientation and densification through a multi-stage composite field activation sintering process (microwave-magnetic field-pressure coupling). A gradient pressure regime is matched to the material's shrinkage characteristics, and combined with asymmetric cooling technology, significantly reducing the probability of sintering defects and ensuring the uniformity of performance in large-size products.

[0025] 4. This invention utilizes a quadratic function gradient structure design to achieve continuously varying performance characteristics from the material surface to the core. The high-hardness surface layer provides excellent wear resistance, the intermediate transition layer buffers stress impacts, and the high-toughness core matrix bears the main load. This structural design enables the material to maintain performance stability under complex working conditions.

[0026] 5. This invention, through the synergistic optimization of composite field activation and gradient process, reduces sintering temperature while shortening holding time. The dynamic pressure loading system reduces ineffective energy loss, and the catalytic effect of rare earth elements accelerates the densification process. Overall, the energy consumption of the preparation is significantly lower than that of traditional methods, aligning with the trend of green manufacturing. Attached Figure Description

[0027] Figure 1 This is a schematic diagram of the preparation method of the present invention. Detailed Implementation

[0028] The following is in conjunction with the appendix Figure 1 : Example

[0029] Raw material ratio (parts by mass): 77.5 parts WC, 10 parts TiC, 6.5 parts TaC, 4 parts ZrC, 2 parts HfC, and 1.0 part rare earth oxides (Y2O3:La2O3:Sc2O3=7:3:0.5). The preparation process is as follows: S1. Preparation of core-shell structure: Pulsed chemical vapor deposition was employed at a deposition temperature of 850℃, an H2 / CH4 volume ratio of 9:1, and 4 pulse cycles. A TiC shell with a thickness of 20 nm and a surface roughness of Ra = 12 nm was obtained. S2, Gradient layer construction: Laser parameters: wavelength 1070nm, power 200W, scanning speed 1000mm / s; Single-layer thickness 50μm, interlayer rotation angle 68°, oxygen content ≤30ppm; Gradient distribution function: for TaC surface content of 100%, f(x) = −0.04x 2 +0.2x decreases to 0%, while ZrC increases to 80% accordingly; S3, Synergistic Activation Treatment: Microwave frequency 2.45 GHz, magnetic field frequency 50 kHz, intensity 0.5 T; Processing time 30 min, temperature gradient 20℃ / mm; S4, Multi-field Coupled Sintering: Pressure loading: Stage 1 (5 MPa / min) → Stage 2 (1.0 MPa / min) → Stage 3 (0.5 MPa / min); Sintering temperature 1500℃, holding time 20min; S5, Gradient Cooling: Phase 1: 50℃ / min (1600→1000℃); Phase 2: 10℃ / min (1000→300℃); Vacuum degree 2×10 -3 Pa.

[0030] Example 2: Raw material ratio (parts by mass): 75 parts WC, 8 parts TiC, 5 parts TaC, 3 parts ZrC, 1.5 parts HfC, and 0.8 parts rare earth oxides (Y2O3:La2O3:Sc2O3 = 6.5:2.5:0.4). The preparation process is as follows: S1. Preparation of core-shell structure: Deposition temperature 830℃, H2 / CH4 volume ratio 8.5:1, pulse cycle 3 times; The TiC shell thickness is 18 nm, and Ra = 15 nm. S2, Gradient layer construction: Laser wavelength 1060nm, power 180W, scanning speed 800mm / s; Single-layer thickness 45μm, interlayer rotation angle 66°, oxygen content 50ppm; Gradient function: f(x) = −0.05x for TaC surface content of 100%. 2 Decrease by +0.25x; S3, Synergistic Activation: Microwave 2.4 GHz, magnetic field 48 kHz, intensity 0.45 T; Processing time 28 min, temperature gradient 25℃ / mm; S4. Sintering process: Pressure loading rate: 4.5→0.8→0.4MPa / min; Sintering temperature 1450℃, holding time 18min; S5, Cooling Control: Stage 1: Cooling at 45℃ / min; Stage 2: Cooling at 8℃ / min; Vacuum degree 5×10 -3 Pa.

[0031] Example 3: Raw material ratio (parts by mass): 80 parts WC, 12 parts TiC, 8 parts TaC, 5 parts ZrC, 2.5 parts HfC, and 1.2 parts rare earth oxides (Y2O3:La2O3:Sc2O3 = 7.5:3.5:0.6); The preparation process is as follows: S1. Preparation of core-shell structure: Deposition temperature 870℃, H2 / CH4 volume ratio 9.5:1, pulse cycle 5 times; TiC shell 22nm, Ra=8nm; S2, Gradient layer construction: Laser wavelength 1080nm, power 220W, scanning speed 1200mm / s; Single-layer thickness 55μm, interlayer rotation angle 70°, oxygen content 10ppm; Gradient function: TaC with respect to f(x) = −0.03x 2 Decrease by +0.15x; S3, Synergistic Activation: Microwave 2.5GHz, magnetic field 52kHz, intensity 0.55T; Processing time: 32 min; Temperature gradient: 15℃ / mm; S4. Sintering process: Pressure loading rate: 5.5→1.2→0.6MPa / min; Sintering temperature 1550℃, holding time 22min; S5, Cooling Control: Stage 1: Cooling at 55℃ / min; Phase 2: Cooling at 12℃ / min; Vacuum degree 1×10 -3 Pa.

[0032] Comparative Example 1: Compared with Example 1, the difference is that uncoated WC powder and TiC powder are directly mechanically mixed, and the core-shell structure preparation in step S1 is omitted. The remaining process parameters are the same as in Example 1.

[0033] Comparative Example 2: Compared with Example 1, the difference is that the raw material ratio is adjusted to 82 parts WC, 13 parts TiC, 9 parts TaC, 6 parts ZrC, 3 parts HfC, and 1.3 parts rare earth oxides, while the remaining process parameters are the same as in Example 1.

[0034] Comparative Example 3: Compared with Example 1, the difference is that only Y2O3 is used as the rare earth oxide, the molar ratio is adjusted to Y2O3:La2O3:Sc2O3=10:0:0, the total addition amount is kept at 1.0 part, and the other process parameters are the same as those in Example 1.

[0035] Comparative Example 4: Compared with Example 1, the difference is that in step S2, equal amounts of mixed powder are directly pressed, and the gradient layer construction process is not implemented. The remaining process parameters are the same as in Example 1.

[0036] Comparative Example 5: Compared with Example 1, the difference is that step S3, the co-activation treatment, is omitted, and step S4 is changed to unidirectional pressure sintering (constant pressure 50MPa, rate 5MPa / min). The remaining process parameters are the same as in Example 1.

[0037] Comparative Example 6: Compared with Example 1, the difference is that step S5 uses a single-rate cooling (20°C / min from 1600°C to room temperature) and does not implement gradient cooling. The remaining process parameters are the same as in Example 1.

[0038] Comparative Example 7: Compared with Example 1, the difference is that in step S1, the deposition parameters (H2 / CH4=7:1, temperature 900℃) are adjusted so that the TiC shell layer and the rare earth oxide layer are in direct contact, and no 2-3nm transition interface layer is formed. The other process parameters are the same as in Example 1.

[0039] Comparative Example 8: Compared with Example 1, the difference is that the TaC content in step S2 decreases linearly (f(x)=−kx+b) rather than as a quadratic function distribution, while the other process parameters are the same as in Example 1.

[0040] Test Example 1: Comparison Experiment of Mechanical Properties Experimental equipment: Vickers hardness tester (Wilson-Wolpert-432SVD, load 30kg); Universal testing machine (Instron-5985, equipped with three-point bending fixture); Surface roughness tester (Mitutoyo-SJ-410); Sample preparation: The sintered bodies of Example 1 and Comparative Examples 1 / 4 / 5 were processed into standard samples: Hardness test: 10×10×5mm polished sample (Ra≤0.1μm); Bending strength: 3×4×40mm long strip specimen (span 30mm); Fracture toughness: Pre-notched specimen (notch depth 2 mm, tip radius 15 μm); The experimental steps are as follows: Vickers hardness test: 1. Select 5 test points on each sample surface, with a spacing ≥ 3 times the diagonal of the indentation; 2. The loading holding time is 15 seconds, and the diagonal length of the indentation is measured after unloading; 3. After removing outliers (deviation > 10%), take the average value; Bending strength test: 1. Loading rate 0.5 mm / min, record load-displacement curves; 2. After fracture, measure the location of the fracture surface to ensure that the fracture occurs in the middle 1 / 3 of the span; 3. The number of valid samples is ≥5; Fracture toughness test: 1. The single-sided notched beam method (SENB) is adopted; 2. Loading rate 0.05 mm / min, record the critical load Pc; 3. Calculate the KIC value according to ASTM-E399 standard. The experimental data are shown in Table 1: Table 1: Comparison Test Table of Mechanical Properties Sample Vickers hardness (HV30) Flexural strength (MPa) <![CDATA[Fracture toughness (MPa·m 1 / 2 )]]> Example 1 2453±87 1896±112 14.2±0.8 Comparative Example 1 1987±134 1423±98 9.6±0.6 Comparative Example 4 2175±121 1632±105 11.3±0.7 Comparative Example 5 2289±95 1754±89 12.1±0.9 Experiment Summary: The superior hardness performance of Example 1 is attributed to the multi-level interface design of the core-shell structure. The strong bonding interface formed by the WC core and TiC shell through lattice coherence effectively hinders dislocation slip. Simultaneously, the transient liquid phase generated during the sintering process of the rare earth oxide coating optimizes the interparticle bonding state. In Comparative Example 1, due to the mechanical mixing method, micropore defects exist at the hard phase interface, leading to stress concentration during indentation testing and a significant decrease in hardness value.

[0041] The difference in flexural strength highlights the synergistic strengthening mechanism of the gradient structure and the composite field process. In Example 1, the quadratic gradient distribution of TaC forms a progressive stress buffer layer, inducing crack bifurcation and propagation under bending loads, thus lengthening the fracture path. In contrast, the homogeneous structure in Comparative Example 4 lacks this gradient transition, and the crack rapidly penetrates the cross-section in a straight line, resulting in a strength decrease of approximately 17%. The non-thermal effect generated by the activation of the composite field promotes atomic diffusion, thereby enhancing the grain boundary bonding strength in Example 1 and causing the fracture surface to exhibit more dimple morphology.

[0042] Fracture toughness data confirm the microstructure regulation effect of multi-stage pressure loading. In Example 1, during the second-stage pressure loading (0.8-1.2 MPa / min), the formation of an interlaced grain boundary network by controlling the plastic flow rate resulted in crack propagation requiring more energy. Comparative Example 5, using a constant pressure rate, led to excessive grain orientation growth and the formation of weak bonding interfaces, resulting in a toughness value approximately 15% lower than that of Example 1. The residual compressive stress field generated by the gradient cooling regime further suppressed surface crack initiation.

[0043] Test Example 2: High Temperature Performance Comparison Experiment Experimental equipment: High-temperature hardness tester (Zwick / Roell-ZHU2.5, maximum temperature 1200℃); Box-type resistance furnace (Nabertherm-LHT-04 / 17). Precision balance (Mettler-Toledo-XS205, accuracy 0.01mg); Sample preparation: The sintered bodies of Example 1 and Comparative Examples 2 / 7 / 8 were processed into: High-temperature hardness test specimen: 8×8×4mm, surface polished to Ra≤0.2μm; Oxidation sample: φ10×2mm round disc, double-sided precision ground; Thermal shock test specimen: 15×15×3mm square block, with edge chamfer R0.5mm; The experimental steps are as follows: High-temperature hardness test: 1. In a vacuum environment (10 -2 Pa) is heated to 1000℃ at a rate of 10℃ / min; 2. After maintaining the temperature for 30 minutes, apply a 10kg load and hold for 20 seconds; 3. Test 3 effective points for each sample, with a spacing of ≥2mm; Oxidative weight gain test: 1. Weigh the sample after cleaning and drying (initial mass m0). 2. Maintain a temperature of 1000℃ for 50 hours in a static air environment; 3. After cooling, ultrasonic cleaning is performed to remove the loose oxide layer, and the final mass m1 is measured; 4. Calculate the weight gain per unit area Δm = (m1 - m0) / A; Thermal shock cycling test: 1. Heat the sample to 1000℃ and hold for 15 minutes; 2. Quickly transfer to 20℃ deionized water for quenching; 3. Visually inspect for surface cracks after every 5 cycles, and monitor cracking signals using an acoustic emission instrument; The experimental data are shown in Table 2: Table 2: High Temperature Performance Comparison Test Table Sample High-temperature hardness (HV10) <![CDATA[Oxidation weight gain (mg / cm 2 )]]> thermal shock cycles Example 1 1276±68 3.2±0.4 28 Comparative Example 2 1055±92 5.8±0.7 17 Comparative Example 7 1189±77 4.1±0.5 23 Comparative Example 8 1124±84 4.9±0.6 19 Experiment Summary: The superior high-temperature hardness of Example 1 stems from the surface enrichment effect of TaC in the gradient structure. Under test conditions of 1000℃, TaC preferentially oxidizes to form a dense Ta2O5 protective film (melting point 1872℃), effectively blocking oxygen diffusion inward. In Comparative Example 2, the TaC content exceeded the reasonable range, resulting in an excessively thick gradient layer that caused thermal expansion mismatch, leading to microcracks at high temperatures and a decrease in hardness value of approximately 17%. The presence of the transition interface layer (Example 1) alleviated the thermal stress between the TiC shell and rare earth oxides, while the directly contacting heterogeneous interface in Comparative Example 7 became a weak zone for high-temperature softening.

[0044] Oxidation weight gain data validated the advantages of the quadratic function gradient distribution. In Example 1, TaC decreased nonlinearly, forming a high concentration gradient in the near-surface region (dC / dx ≈ −0.08x), resulting in a continuous transition structure in the oxide film. In Comparative Example 8, the linear distribution (dC / dx = constant) led to a compositional abrupt change at the oxide film / substrate interface, generating internal stress during oxidation that caused the protective film to peel off, increasing the weight gain rate by 53%. The addition of Sc2O3 promoted oxide film grain refinement (average grain size <500nm), enhancing its resistance to peeling.

[0045] The difference in thermal shock cycle life reflects the synergistic effect of gradient design and interface engineering. In Example 1, the increasing distribution of ZrC creates a gradual transition in the coefficient of thermal expansion (surface 8.2 × 10⁻⁶). -6 / ℃→Internal 7.1×10 -6 / ℃), reducing the thermal stress amplitude. In Comparative Example 7, due to the lack of a transition interface layer, microcracks initiated and rapidly propagated at the interface during thermal shock. The constant distribution of HfC (20%) promotes uniform heat dissipation through high thermal conductivity (20 W / m·K), avoiding failure caused by local overheating.

[0046] Test Example 3: Structural Stability Comparison Experiment Experimental equipment: X-ray stress analyzer (Proto-LXRD, Cr-Kα radiation); Electronic density meter (Quantachrome-Ultrapyc-1200e); Thermal expansion meter (Netzsch-DIL-402C); Sample preparation: The sintered bodies of Example 1 and Comparative Example 3 / 6 were processed into: Residual stress specimen: 20×20×5mm flat plate, surface electropolished; Density sample: φ6×8mm cylinder, end face parallelism ≤0.01mm; Thermal expansion specimen: 25×5×5mm square bar, surface roughness Ra≤0.4μm; The experimental steps are as follows: Residual stress test: 1. Using sin 2 ψ method, select (311) diffraction plane (2θ≈136°); 2. Establish a 5×5 grid of test points on the sample surface, with a spacing of 2 mm; 3. Calculate the surface normal stress σ and stress gradient dσ / dz; Density measurement: 1. Volume measurement using helium displacement method, accuracy ±0.01%; 2. Each sample was measured three times, and the average value was taken. 3. Calculate the relative density: (Measured density / Theoretical density) × 100%; Thermal expansion coefficient test: 1. Heating rate: 5℃ / min; Temperature range: 25-1000℃; 2. Record the curve of ΔL / L0 as a function of temperature; 3. Calculate the average linear expansion coefficient (25-800℃ range); The experimental data are shown in Table 3: Table 3: Structural Stability Test Table Sample Surface residual stress (MPa) Relative density (%) <![CDATA[Coefficient of linear thermal expansion (×10 -6 / °C)]]> Example 1 -218±35 99.5±0.2 7.3±0.1 Comparative Example 3 +156±42 98.1±0.3 7.9±0.2 Comparative Example 6 -85±28 99.0±0.2 7.6±0.1 Experiment Summary: The surface compressive stress distribution in Example 1 confirms the effectiveness of the gradient cooling regime. The rapid cooling stage (50°C / min) induces a martensitic transformation on the surface, while the subsequent slow cooling stage (10°C / min) allows for gradual stress release in the core, resulting in a residual compressive stress layer of -218 MPa. Comparative Example 6, using a single cooling rate, suffers from core shrinkage constrained by the surface layer, leading to uneven stress distribution and a 61% reduction in compressive stress. The addition of La2O3 promotes grain boundary slip, effectively releasing localized stress in the later stages of sintering and avoiding the tensile stress state observed in Comparative Example 3.

[0047] The density difference reflects the promoting effect of rare earth synergy on densification. In Example 1, Y... 3+ Grain boundary segregation reduces diffusion activation energy (by approximately 15%), Sc 3+ The liquid phase wetting effect fills the micropores, achieving a relative density of 99.5%. Comparative Example 3 contains only Y₂O₃ and lacks Sc. 3+ Liquid-phase assisted densification resulted in the presence of isolated pores (approximately 0.5 μm) at grain boundaries, leading to a 1.4% decrease in density. A gradient sintering pressure regime further reduced closed-cell formation by matching the shrinkage characteristics at different stages.

[0048] The variation in the coefficient of thermal expansion validates the thermal matching design of the gradient structure. In Example 1, the continuous gradient distribution of ZrC (80% on the surface → 20% inside) makes the overall coefficient of thermal expansion of the material approach that of the matrix (7.3 × 10⁻⁶). -6 / ℃), a 7.6% reduction compared to Comparative Example 3. The uniform distribution of HfC (2%) stabilizes the grain boundary structure at high temperatures through its high melting point (3928℃), suppressing abrupt expansion caused by abnormal grain growth. The uniform heating effect generated by microwave-magnetic field synergistic activation ensures that the gradient composition distribution accurately achieves the design target.

[0049] Test Example 4: Comparison Experiment of Process Characteristics Experimental equipment: Laser displacement sensor (Keyence-LK-G5000, accuracy ±0.5μm); Thermal imager (FLIR-A655sc, sampling rate 100Hz); Power analyzer (Yokogawa-WT1800, accuracy 0.1%). Sample preparation: Using the original sintered green bodies from Example 1 and Comparative Examples 5 / 6: Densification sample: φ20×10mm cylindrical blank (unsintered); Thermal stress simulation specimen: 50×50×5mm flat blank; Energy consumption test sample: mixed powder from the same batch (200g / set); The experimental steps are as follows: Sintering densification monitoring: 1. Install a laser displacement sensor in the observation window of the sintering furnace; 2. Record the changes in billet height in real time (sampling interval 10s); 3. Calculate the densification rate v = Δh / (h0·Δt); Thermal stress simulation: 1. Establish a three-dimensional finite element model of the billet (mesh size 0.5mm); 2. Import measured temperature field data (acquired by thermal imager); 3. Calculate the distribution of maximum principal stress during the cooling process; Energy consumption measurement: 1. Monitor the input power of the sintering furnace throughout the entire process; 2. Calculate the energy consumption per unit mass, E = P·t / m (kWh / kg); 3. Repeat the experiment three times and take the average value; The experimental data are shown in Table 4: Table 4: Comparison Test Table of Process Characteristics Sample Maximum densification rate (% / min) Peak thermal stress (MPa) Energy consumption per unit (kWh / kg) Example 1 0.85±0.07 218±15 4.2±0.3 Comparative Example 5 0.63±0.09 327±22 5.8±0.4 Comparative Example 6 0.72±0.08 285±18 4.9±0.2 Experiment Summary: The increased densification rate in Example 1 stems from the synergistic activation effect of microwaves and magnetic fields. The non-thermal effects of microwaves promote grain boundary diffusion (increasing the diffusion coefficient by approximately 40%), while the eddy currents induced by the alternating magnetic field generate localized Joule heating (a temperature rise of approximately 50°C), reducing the sliding frictional resistance between powder particles. Comparative Example 5, using conventional sintering with a higher activation energy barrier, primarily relies on thermal diffusion for densification, resulting in a 26% decrease in the densification rate. Multi-stage pressure loading (0.2→0.8→1.2 MPa / min) matches the deformation characteristics of different sintering stages, avoiding the insufficient particle rearrangement caused by a single loading rate in Comparative Example 6.

[0050] The reduction in thermal stress validated the rationality of the gradient cooling regime. In Example 1, rapid cooling (50°C / min) at the high-temperature stage formed a surface compressive stress layer, while slow cooling (10°C / min) at the intermediate temperature stage allowed for core stress release, resulting in a 24% reduction in peak thermal stress compared to Comparative Example 6. The liquid-phase wetting effect of rare earth oxides filled the grain boundary micropores, increasing the material's high-temperature strength (by approximately 15%) and enhancing its resistance to thermal deformation. In Comparative Example 5, due to the lack of gradient cooling, core shrinkage was hindered, leading to tensile stress concentration and a 50% increase in the maximum principal stress value.

[0051] The difference in energy consumption reflects the comprehensive benefits of process optimization. Example 1 reduced the sintering temperature (by approximately 80°C) by activating the composite field, shortened the holding time by 30%, and reduced unit energy consumption by 28%. The gradient pressure regime reduced ineffective plastic deformation work, and compared with the constant pressure of Comparative Example 5, the energy utilization rate was improved by 22%. The uniform distribution of HfC (2%) improved the thermal conductivity of the material (by approximately 18%), promoted the homogenization of the temperature field during sintering, and reduced heat loss.

[0052] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A binderless cemented carbide, characterized in that, The components include the following parts by mass: WC: 75-80 servings; TiC: 8-12 parts; TaC: 5-8 servings; ZrC: 3-5 parts; HfC: 1.5-2.5 parts; Rare earth oxides: 0.8-1.2 parts.

2. The binderless cemented carbide according to claim 1, characterized in that, The WC and TiC form a core-shell structure. The WC core has a particle size of 180-220 nm, and the TiC shell has a thickness of 18-22 nm. The surface of the core-shell structure is covered with a rare earth oxide layer with a thickness of 10-15 nm. A 2-3 nm transition interface layer is formed between the TiC shell and the rare earth oxide layer.

3. The binderless cemented carbide according to claim 1, characterized in that, The rare earth oxides include Y2O3, La2O3 and Sc2O3. In the mixture of Y2O3, La2O3 and Sc2O3, the molar ratio of the three is 6.5-7.5:2.5-3.5:0.4-0.6, and the rare earth oxides exist on the surface of the core-shell structure in the form of a continuous coating layer with a coating layer thickness of 10-15 nm. The particle size distribution of Y2O3, La2O3 and Sc2O3 in the coating layer is D90≤200 nm.

4. The binderless cemented carbide according to claim 1, characterized in that, The TaC, ZrC and HfC form a continuous gradient distribution along the thickness direction. The TaC content decreases from the material surface to the interior according to a quadratic function law, with the surface content decreasing from 100% to 0%. The ZrC content increases from 0% to 80%, and the HfC content remains constant at 20%. The total thickness of the gradient layer is 15-25% of the material thickness.

5. A method for preparing a binderless cemented carbide, applied to the binderless cemented carbide according to any one of claims 1-4, characterized in that, Includes the following steps: S1. Core-shell structure preparation: A continuous titanium carbide shell layer is formed on the surface of tungsten carbide powder by vapor deposition to obtain a core-shell structured composite powder. S2. Gradient layer construction: The core-shell structured composite powder and gradient phase powder are stacked layer by layer using high-energy beam melting technology to form a composition gradient distribution that changes continuously along the thickness direction; S3. Synergistic activation treatment: Under the synergistic effect of the combined field of microwave and alternating magnetic field, the gradient structure blank is activated across scales. S4. Multi-field coupling sintering: Plasma activation sintering with multi-stage pressure loading is applied to the activated green body to achieve material densification and microstructure stabilization. S5. Gradient cooling treatment: The sintered body is subjected to stress regulation by a staged variable rate cooling method to release the residual stress gradient.

6. The method for preparing a binderless cemented carbide according to claim 5, characterized in that, In step S1: The vapor deposition method is pulsed chemical vapor deposition, with the deposition temperature controlled at 830-870℃, the volume ratio of hydrogen to methane in the carrier gas being 8.5-9.5:1, and the pulse period being 3-5 times. After deposition, the titanium carbide shell thickness in the core-shell structured composite powder is 18-22nm, and the surface roughness Ra of the shell is ≤15nm.

7. The method for preparing a binderless cemented carbide according to claim 5, characterized in that, In step S2: The high-energy beam melting technology uses a laser beam with a wavelength of 1060-1080nm, a laser power of 180-220W, a scanning speed of 800-1200mm / s, a single-layer stacking thickness of 45-55μm, an interlayer rotation angle of 66°-70°, and an oxygen content of ≤50ppm in the protective atmosphere.

8. The method for preparing a binderless cemented carbide according to claim 5, characterized in that, In step S3: The microwave frequency is 2.4-2.5 GHz, the alternating magnetic field frequency is 48-52 kHz, the magnetic field strength is 0.45-0.55 T, the processing time is 28-32 min, and the temperature gradient of the billet is controlled within 15-25 ℃ / mm during the processing.

9. The method for preparing a binderless cemented carbide according to claim 5, characterized in that, In step S4: The multi-stage pressure loading includes three stages: The first stage involves loading at a rate of 4.5-5.5 MPa / min within a pressure range of 0-30 MPa. The second stage involves loading at a rate of 0.8-1.2 MPa / min within the range of 30-45 MPa; The third stage involves loading at a rate of 0.4-0.6 MPa / min within the range of 45-50 MPa; The sintering temperature is maintained at 1450-1550℃, and the holding time is 18-22 minutes.

10. The method for preparing a binderless cemented carbide according to claim 5, characterized in that, In step S5: The gradient cooling includes a first stage of cooling from 1600℃ to 1000℃ at a rate of 45-55℃ / min, a second stage of cooling from 1000℃ to 300℃ at a rate of 8-12℃ / min, and a third stage of natural cooling to room temperature, maintaining a vacuum degree ≤5×10⁻⁵ throughout the cooling process. -3 Pa.

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