Multi-element composite alloy additive for high hardenability plastic mold steel and application thereof

Abstract

The application relates to an alloy additive for plastic mold steel, discloses a multi-element composite alloy additive for high quenching plastic mold steel and application of the multi-element composite alloy additive, and aims to solve the problems that the porosity, the strength and the magnetic permeability of the porous mold steel are mutually restricted and the pores are easily blocked by polymer residues. The additive is a composite powder, the soft magnetic alloy nanoparticles are taken as the core, the outer surface of the soft magnetic alloy nanoparticles is coated with a transient liquid phase alloy shell layer with a liquidus temperature of 850 DEG C to 920 DEG C and the highest mass fraction of metal elements with a melting point not lower than 1400 DEG C, and the outer surface of the shell layer is attached with two-dimensional layered ceramic nanosheets which are not wetted by the alloy melt. During sintering, the melt carries the soft magnetic phase to the sintering neck and is isothermally solidified, and the nanosheets are repulsively migrated to the pore wall to form a low surface energy isolation layer. The additive can be used for preparing an exhaust insert of an injection mold.

Description

Technical Field

[0001] This invention relates to alloy additives for plastic mold steel, and particularly to a multi-component composite alloy additive for high hardenability plastic mold steel and its application. Background Technology

[0002] In precision injection molding, gas in deep cavities and structural dead corners is difficult to escape at the moment of mold closing. The resulting localized trapped gas can cause scorching, insufficient glue, air bubbles, and localized stress concentration in the molded part. To eliminate this defect, the industry practice is to use powder metallurgy to prepare porous, breathable mold steel and process it into venting inserts to be installed in poorly vented areas. The typical route uses martensitic stainless steel powder or chromium-containing hot work die steel powder as the skeleton material, and physically mixes 10% to 15% by mass of organic pore-forming agents, such as polymethyl methacrylate microspheres, urea or ammonium bicarbonate. The green body is obtained by cold isostatic pressing or molding, and then the pore-forming agent is removed by slow heating under vacuum or protective atmosphere. Subsequently, it is held at a conventional sintering temperature range of 1250°C to 1350°C, so that the metal particles form sintering necks at the contact points by relying on surface diffusion and bulk diffusion. Finally, a connected pore network with a total porosity of 15% to 25% and a pore size of several micrometers to more than ten micrometers is obtained. The cavity gas is discharged through the pore network under the action of pressure difference according to Darcy's law.

[0003] With the widespread application of rapid heating and cooling injection molding technology in high-gloss paint-free interior parts, thin-walled communication structural parts, and precision optical components, the mold surface needs to be heated within seconds using medium- and high-frequency electromagnetic induction to eliminate weld lines. This requires venting inserts to not only have venting capabilities but also be induction-heated. However, porous bodies formed by solid-state sintering cannot meet this requirement. Pore-forming agents artificially widen the particle spacing, and to prevent the pores from shrinking and closing at high temperatures, the sintering temperature and holding time must be limited. The resulting sintered neck cross-section is usually less than a few percent of the particle cross-section, becoming a dual bottleneck in mechanics and magnetism. Under load, stress is highly concentrated here, and the tensile strength of the material hovers below 500 MPa for a long time. During magnetization, the magnetic flux is forced to be transmitted entirely through these narrow necks. Due to their small cross-section, the necks saturate first, and the effective permeability of the porous body is thus dominated by the proportion of the neck cross-section, which is an order of magnitude lower than that of dense steel of the same composition. The demagnetizing field generated by the pores themselves is only a secondary additional factor. Since the surface power density of induction heating depends on the material's permeability and resistivity, excessively low permeability directly leads to insufficient power deposition per unit excitation current, making it difficult to achieve the required heating rate on the mold surface. If density is increased to improve strength and permeability, gas permeability decreases sharply according to the cubic relationship of porosity, resulting in the loss of venting function. This creates an irreconcilable constraint between porosity and mechanical and electromagnetic properties.

[0004] Another long-standing defect stems from the surface chemical state of the pore walls. The pore walls left after the pore-forming agent escapes are essentially free surfaces of unpassivated metal crystals. These surfaces are enriched with active iron and chromium atoms, possessing numerous unpaired dangling bonds, resulting in high surface free energy and chemical reactivity. When processing amorphous special engineering plastics like polyetherimide, whose glass transition temperature is approximately 217°C and melt processing temperature ranges from 340°C to 415°C, oligomers, small molecule pyrolysis products, and mold release agent residues generated during melt decomposition readily undergo chemical adsorption at these high-energy sites and further cross-link and carbonize. This leads to continuous scaling and accumulation within pores only a few micrometers wide, causing irreversible degradation of air permeability after tens of thousands of cycles. Coating processes such as chemical vapor deposition and physical vapor deposition, which have been attempted to address this problem, are limited by the diffusion depth of the gas phase in the three-dimensional tortuous channels and cannot penetrate deep into the porous body. Organic passivation layers obtained by impregnation with fluorinated surfactants are rapidly degraded and detached under repeated scouring by high-temperature melts and high-pressure gas flows, and similarly cannot provide long-term protection.

[0005] Therefore, the technical problems to be solved by the present invention are as follows: First, without reducing the pore connectivity and gas permeability of porous plastic mold steel, the macroscopic tensile strength and magnetic permeability are simultaneously improved, overcoming the mutual constraints between porosity and mechanical and electromagnetic properties; Second, the dependence on secondary coating post-treatment is eliminated, and a low surface energy protective layer is formed in situ on the inner wall of the micron-level interconnected channels during the sintering process, suppressing the pore blockage and air permeability reduction caused by the adsorption and accumulation of polymer residues during long-term high-temperature injection molding service. Summary of the Invention

[0006] To address the aforementioned technical problems, this invention provides a multi-component composite alloy additive for high hardenability plastic mold steel and its application. The core of this invention lies in pre-assembling the magnetic functional phase, metallurgical bonding phase, and interface modification phase required for modification into a composite powder with a defined spatial arrangement at the single-particle scale. This transforms the sintering process from passive solid-state atomic diffusion to directional material transport and self-assembly driven by capillary force and interfacial free energy.

[0007] According to one aspect of the present invention, a multi-component composite alloy additive for high hardenability plastic mold steel is provided. The multi-component composite alloy additive is a composite powder. Each particle of the composite powder comprises soft magnetic alloy nanoparticles, a transient liquid phase alloy shell coating the outer surface of the soft magnetic alloy nanoparticles, and two-dimensional layered ceramic nanosheets attached to the outer surface of the transient liquid phase alloy shell. The transient liquid phase alloy contains at least one melting point reducing element selected from phosphorus, boron, and silicon. The liquidus temperature of the transient liquid phase alloy is 850°C to 920°C, and the melting point of the metal element with the highest mass fraction in the transient liquid phase alloy is not lower than 1400°C. The wetting contact angle of the melt of the transient liquid phase alloy on a dense substrate of the same material as the two-dimensional layered ceramic nanosheets is greater than 90°. The wetting contact angle is a value measured by the seat drop method under the conditions of temperature 950°C and vacuum degree not higher than 10⁻³ Pa.

[0008] The three components mentioned above play irreplaceable roles in the system and are triggered sequentially in time. Soft magnetic alloy nanoparticles provide high saturation magnetization, but they lack reliable metallurgical bonding with the martensitic matrix. If added alone, they can only be randomly distributed as isolated islands in the powder aggregate, failing to form magnetic flux pathways at the sintering neck and introducing soft, weak regions at the particle contact points. The transient liquid-phase alloy shell functions as both a transport carrier and an active solder in the system. It melts first during heating, and the soft magnetic alloy nanoparticles it coats then transform into dispersed particles suspended in the melt. As the melt, driven by capillary force, converges towards the wedge-shaped gap with the smallest radius of curvature at the matrix particle contact point, the nanoparticles are carried there by viscous drag force. Thus, the magnetic functional phase is precisely deployed at the sintering neck, where magnetic resistance and stress concentration are most severe. The liquidus temperature of the alloy is limited to 850°C to 920°C to ensure that the melt is fully formed in the temperature range before the matrix powder undergoes perceptible solid-phase sintering shrinkage, thus providing sufficient time and space for redistribution. When the liquidus temperature is below 850°C, the melt is formed before degreasing is completely finished, and residual organic matter decomposition gases will form pores in the liquid film and disrupt its continuity. When the liquidus temperature is above 920°C, the holding temperature required for complete melting of the shell will approach the temperature range where the matrix powder begins to shrink, and the porous framework will begin to densify before the melt completes its flow.

[0009] The requirement that the melting point of the metallic element with the highest mass fraction in the transient liquid phase alloy be no less than 1400℃ is a prerequisite for achieving the key mechanism of isothermal solidification. The essence of transient liquid phase bonding lies in the fact that during the holding process, the melting point-reducing element diffuses and flows away towards the steel matrix along the concentration gradient, while the iron in the matrix dissolves into the bonding zone in the opposite direction. The composition of the bonding zone continuously shifts towards the main metal side of the alloy, and its melting point increases accordingly. Only when the melting point of the main metal itself is significantly higher than the holding temperature can the composition shift cause the melting point of the bonding zone to exceed the ambient temperature and transform from liquid to solid under isothermal conditions. If copper, with a melting point of only 1085℃, is chosen as the main metal, then no matter how much the melting point depletion element diffuses and is lost, the upper limit of the melting point of the residual liquid phase can only approach the melting point of copper. In the holding temperature range above 1150℃, the liquid phase will always remain liquid, and isothermal solidification will not occur. At this time, the residual copper-based liquid phase will not only fail to form a solid solution-strengthened bridging layer, but will also penetrate along the austenite grain boundaries of the matrix, causing the liquid metal to become brittle, which will actually reduce the strength and hardness of the sintered body.

[0010] The requirement that the wetting contact angle between the two-dimensional layered ceramic nanosheets and the transient liquid-phase alloy melt be greater than 90° is a necessary condition for achieving in-situ modification of the pore walls. To minimize the total interfacial free energy, the system spontaneously repels the non-wetting nanosheets to the three-phase interface of solid, liquid, and pores during melt contraction and convergence, allowing them to adhere extensively to the inner surface of the pores after solidification. If the two wet each other, the nanosheets will be engulfed by the melt and ultimately trapped inside the metal, leaving the pore walls in a bare, highly active metallic state, rendering any anti-clogging effect meaningless. Since the droplet method can only be implemented on a dense, flat substrate, this feature uses a dense substrate of the same material as the nanosheets as the test subject, and specifies a test temperature of 950°C. This temperature is higher than the upper limit of the liquidus temperature defined in this invention, ensuring the presence of a stable melt in all covered alloy systems.

[0011] Preferably, based on the total mass of the multi-component composite alloy additive, the mass fraction of the soft magnetic alloy nanoparticles is 12% to 25%, the mass fraction of the two-dimensional layered ceramic nanosheets is 1% to 5%, and the remainder is the transient liquid phase alloy and unavoidable impurities. This ratio range essentially constrains the solid volume fraction of the suspended system after the shell melts. Based on the density of each phase, when the mass fraction of the soft magnetic alloy nanoparticles is 25%, the volume fraction of suspended particles in the melt is approximately 26%. According to the Krieger-Dougherty relationship, taking the maximum bulk volume fraction of 0.6, its relative apparent viscosity is approximately 2.3 times the intrinsic viscosity of the melt, and capillary-driven directional flow can still proceed smoothly. If the mass fraction exceeds 25%, the volume fraction of the suspended solid phase rapidly approaches the maximum bulk fraction, and the relative viscosity diverges according to this relationship. The system changes from a flowable suspended melt to a semi-solid paste that cannot undergo directional transport. When the mass fraction is below 12%, the total amount of high magnetization material transported to the sintering neck is insufficient, making it difficult to continuously connect the magnetic flux pathways in the neck, thus limiting the improvement in magnetic permeability. When the content of two-dimensional layered ceramic nanosheets is below 1%, the spreading area on the pore walls is insufficient to form a continuous barrier, leaving a considerable proportion of exposed high-energy sites; when it is above 5%, the nanosheets are stacked too thickly on the outer surface of the shell, which hinders the direct contact between the melt and the matrix particles, weakening the metallurgical bonding strength of the sintering neck.

[0012] Preferably, the soft magnetic alloy nanoparticles are selected from at least one of Fe-Co alloy, Fe-Co-V alloy, and Fe-Ni alloy, and the particle size of the soft magnetic alloy nanoparticles is 50 nm to 200 nm. When the particle size is less than 50 nm, the particle specific surface area is too large and the surface energy is too high, which easily leads to agglomeration and surface oxidation during the pretreatment and heating process of chemical plating, resulting in a decrease in magnetic properties. When the particle size is greater than 200 nm, the viscous drag force on the particles in the melt is insufficient to overcome their own inertia and sedimentation tendency, making it difficult for them to migrate with the melt to the sintering neck, and the targeted enrichment effect is significantly weakened. More preferably, the soft magnetic alloy nanoparticles are Fe-Co-V alloy, with Fe accounting for 50%, Co accounting for 48%, and V accounting for 2% by mass percentage. This ratio is close to the maximum value range of the saturation magnetic induction intensity of the Fe-Co binary system, and the trace amount of vanadium can agglomerate at the interface during the high-temperature stage, delaying the dissolution and coarsening of particles during the heat preservation process.

[0013] Preferably, the transient liquid phase alloy is a Ni-Cu-P alloy, with P accounting for 9% to 12% by mass, Cu accounting for 5% to 15%, and the balance being Ni and unavoidable impurities. Nickel is used as the main metal because its melting point of 1455℃ is much higher than the temperature of the second holding stage, providing thermodynamic space for isothermal solidification; furthermore, nickel and iron are infinitely miscible in the solid state, and there is no brittle intermetallic compound interface between the bridging layer and the martensitic matrix after solidification. Phosphorus is a melting point reducing element, and its content determines the melt formation temperature. When the content is below 9%, the alloy deviates significantly from the eutectic composition, and the liquidus temperature rapidly rises outside the range defined by this invention; when the content is above 12%, brittle phosphides are easily left after cooling, and the total phosphorus content in the sintered body exceeds the conventional control level for powder metallurgy iron-based materials. The role of copper is to further lower the eutectic temperature of the system, improve the spread of the melt on the iron-based surface, and enhance its corrosion resistance after diffusing into the matrix. When the copper content is higher than 15%, the role of the main metal nickel is weakened, and the degree of isothermal solidification decreases. During the holding stage, nickel and copper diffuse into the deeper layers of the matrix lattice along the concentration gradient. Both are elements that reduce the critical cooling rate of steel, which allows the porous body to obtain a uniform martensitic structure along the cross-section under gas quenching conditions, thereby improving hardenability.

[0014] Preferably, the two-dimensional layered ceramic nanosheets are hexagonal boron nitride nanosheets, with a thickness of no more than 5 nm and a lateral diameter of 200 nm to 500 nm. Hexagonal boron nitride possesses a graphite-like layered structure, extremely low intrinsic surface energy, and excellent high-temperature stability in an inert atmosphere. Its basal surface does not contain dangling bonds, making it neither wetted by nickel-based melts nor prone to forming chemical bonds with polymer high-temperature decomposition products. When the thickness exceeds 5 nm, the flexibility of the sheets decreases, making it difficult to adhere and spread on pore walls with complex curvatures, and easily forming bridging-like accumulations in the channels. When the lateral diameter is less than 200 nm, the coverage area of ​​a single sheet is too small, requiring a higher addition amount to achieve continuous coverage. When the diameter is greater than 500 nm, the sheets are prone to curling and agglomeration during melt flow, which can block the fine channels.

[0015] According to another aspect of the present invention, a method for preparing the above-mentioned multi-component composite alloy additive is provided, comprising dispersing soft magnetic alloy nanoparticles after sensitization and activation treatment in a chemical plating solution containing nickel salt, copper salt, and hypophosphite for plating, thereby depositing a transient liquid phase alloy shell layer on the outer surface of the nanoparticles to obtain coated particles; and mixing the coated particles with two-dimensional layered ceramic nanosheets in an inert atmosphere using a low-energy dry method, so that the two-dimensional layered ceramic nanosheets adhere to the outer surface of the coated particles. The use of chemical plating rather than mechanical alloying to construct the shell layer is the fundamental difference between this method and conventional powder composite processes. Nickel-based alloys with phosphorus content close to the eutectic composition are brittle due to the presence of a large number of intermetallic compounds. They lack the plasticity required for repeated stretching and cold welding to coat the hard phase in high-energy ball milling. Mechanical alloying in this ductile-deficient system can only produce a mixed powder dominated by fragmentation rather than a coated structure. Chemical plating, on the other hand, relies on the spontaneous oxidation of hypophosphite on the catalytic surface and the reduction of nickel and copper ions. The deposited layer grows with uniform thickness along the particle contour, resulting in a shell of controllable thickness and uniform composition on the nanoparticle surface. The deposited shell is a phosphorus-containing amorphous or nanocrystalline structure, which crystallizes upon heating and melts as a whole near the eutectic temperature. Two-dimensional layered ceramic nanosheets, relying on their easy interlayer slippage, adhere to and spread on the surface of coated particles through van der Waals interactions in low-energy dry mixing. This process does not require plastic deformation and is therefore not limited by the brittleness of the shell.

[0016] Preferably, the pH value of the electroless plating solution is 4.5 to 5.5, and the temperature is 80°C to 90°C. The mass fraction of the transient liquid phase alloy shell is achieved by controlling the ratio of the total amount of metal ions in the plating solution to the amount of soft magnetic alloy nanoparticles. The rotation speed of the low-energy dry mixing is 150 r / min to 250 r / min, and the net running time is 30 min to 60 min. When the pH value is below 4.5, the reducing power of hypophosphite decreases, the deposition rate is too low, and the phosphorus content of the coating is too high. When the pH value is above 5.5, the stability of the plating solution deteriorates, and spontaneous decomposition easily occurs, generating free nickel powder in the liquid phase. When the temperature is below 80°C, the deposition rate is too low. When the temperature is above 90°C, the uniformity of the coating along the particle surface decreases, accompanied by plating solution decomposition. When the rotation speed of low-energy dry mixing is higher than 250 r / min or the time is longer than 60 min, the collision energy is sufficient to cause the brittle shell to peel off and break, and the coating structure is destroyed; when the speed is lower than 150 r / min or the time is shorter than 30 min, the nanosheets are difficult to fully depolymerize and attach uniformly.

[0017] According to another aspect of the present invention, the application of the above-mentioned multi-component composite alloy additive in the preparation of porous and breathable plastic mold steel is provided. The application includes mixing the multi-component composite alloy additive with plastic mold steel matrix powder, molding, and performing stepped heating sintering. The amount of the multi-component composite alloy additive is 5% to 15% of the mass of the plastic mold steel matrix powder. The stepped heating sintering includes a first holding stage and a second holding stage. The temperature of the first holding stage is higher than the liquidus temperature of the transient liquid phase alloy and lower than the solidus temperature of the plastic mold steel matrix powder. The temperature of the second holding stage is higher than the temperature of the first holding stage and lower than the lower of the solidus temperature of the plastic mold steel matrix powder and the melting point of the metal element with the highest mass fraction in the transient liquid phase alloy. When the amount is less than 5%, the total amount of melt generated is insufficient to form a continuous envelope between the three-dimensional contact points of adjacent matrix particles. The transported soft magnetic phase can only be distributed sporadically in an isolated state, and the magnetic flux path in the neck cannot be continuous. When the amount is more than 15%, the excess melt will cause uncontrollable overflow after filling the contact points and invade the main air gap that is intended to be preserved. The porous body will transform from an open-pore network to a closed-pore sintered body, the gas permeability will drop sharply, and the total amount of demelting elements in the sintered body will be too high. Dividing the sintering process into two temperature-progressive holding stages is a prerequisite for achieving the aforementioned micro-reorganization. The first holding stage provides a sufficient kinetic time window for melt generation, capillary convergence, soft magnetic phase transport, and non-wetting phase repulsion. The second holding stage allows the melting point de-melting elements in the melt to diffuse into the deeper layers of the matrix, and the iron in the matrix to dissolve into the bonding region. After the composition of the bonding region shifts towards the main metal, the melting point rises and exceeds the ambient temperature, completing isothermal solidification under isothermal conditions. The second holding stage is simultaneously constrained by two upper limits: the solidus temperature of the matrix and the melting point of the main metal. The former is used to prevent the matrix powder itself from forming a liquid phase, which would cause the porous framework to collapse, while the latter is used to ensure that the bonding region after composition shift actually falls into the solid-state stability range.

[0018] Preferably, the plastic mold steel matrix powder is martensitic stainless steel powder or hot work die steel powder with a particle size of 15 μm to 45 μm; the molding is cold isostatic pressing, with a pressure of 150 MPa to 250 MPa and a holding time of 5 min; the stepped heating sintering includes a degreasing stage in a hydrogen-containing reducing atmosphere, heating at 1℃ / min to 3℃ / min to 430℃ to 480℃ and holding for 2.5 h to 3.5 h; a first holding stage in a vacuum degree not exceeding 10⁻² Pa, heating at 2℃ / min to 4℃ / min to 930℃ to 960℃ and holding for 1.5 h to 3 h; and a heating stage in a high-purity argon atmosphere with a partial pressure of 200 Pa to 1000 Pa, heating at 4℃ / min to 6℃ / min to 1150℃ to 1170℃ and holding for 2.5 h to 3.5 h. The second heat preservation stage of h; the mass fraction of the melting point reducing element in the porous and breathable plastic mold steel obtained after sintering is 0.3% to 0.8%. When the particle size of the matrix powder is less than 15 μm, the original gap between the particles is too fine, and it is difficult to obtain sufficient permeability even without sintering; when it is greater than 45 μm, the pore size is too large, which is beneficial for venting but will damage the surface processing accuracy of the insert and the appearance of the part. When the cold isostatic pressing pressure is lower than 150 MPa, the green strength is insufficient and it is easy to crack during transportation and degreasing; when it is higher than 250 MPa, the spherical particles undergo significant plastic deformation, the original gap is excessively compressed, and there is insufficient three-dimensional space left for melt flow and final channels. The degreasing stage must be carried out in a hydrogen-containing reducing atmosphere and the temperature must be raised to above 430°C. The reason is that the main chain breakage of the binder and the removal of carbonization products need to be higher than the temperature range for removing side groups. Simple vacuum conditions will promote carbon residue, while the presence of hydrogen can convert residual carbon into methane for discharge, and at the same time reduce the initial oxide film on the powder surface, creating conditions for the subsequent spread of melt. In the first holding stage, when the temperature is below 930℃, the shell layer is not completely melted, and the residual solid skeleton makes the apparent viscosity of the melt too high, resulting in a loss of fluidity. Above 960℃, the interdiffusion between the melt and the matrix intensifies prematurely, and the melt is absorbed and dissipated by the matrix before completing redistribution. At the same time, the matrix powder begins to show perceptible shrinkage. In the second holding stage, when the temperature is below 1150℃, the diffusion flux of the demelting element is low, and isothermal solidification is difficult to complete within the holding time. Above 1170℃, on the one hand, the pores undergo spheroidization and shrinkage driven by surface tension; on the other hand, if boron-containing demelting elements are used, the temperature will approach the eutectic temperature of iron and borides, posing a risk of grain boundary liquefaction. The role of the inert gas partial pressure refill in the second holding stage is to suppress the evaporation loss of high vapor pressure elements such as phosphorus and copper under high temperature and long-term holding. If a high vacuum is maintained, some of the demelting elements will leave the bonding zone through evaporation rather than diffusion, and the composition evolution path of isothermal solidification will not be as expected. At the same time, the nickel and copper content on which the hardenability of the matrix depends is also weakened.The mass fraction of the reducing element in the sintered body is controlled between 0.3% and 0.8%. The lower limit ensures sufficient liquid phase generation in the bonding region, while the upper limit is comparable to the conventional control level of this element in powder metallurgy iron-based materials, thus avoiding excessive segregation at grain boundaries. More preferably, the amount of the multi-component composite alloy additive is 8% to 12% of the mass of the plastic mold steel matrix powder.

[0019] Compared with the prior art, the present invention has the following beneficial effects: Because the composite additive pre-establishes the spatial arrangement relationship between the soft magnetic alloy nanoparticles, the transient liquid phase alloy shell, and the two-dimensional ceramic nanosheets on the outer surface at the single-particle scale, during sintering, the melt, driven by capillary force, converges towards the sintering neck and simultaneously carries the soft magnetic phase via suspension transport. The narrow point contact is widened into a solid solution-strengthened bridging band formed by isothermal solidification, allowing magnetic flux and stress to be transmitted through the thickened neck enriched with highly magnetized material. Within the composition and process scope defined in the claims, the total porosity of the porous body is 13.2% to 22.8%, and the gas permeability is 1.4 × 10⁻¹. 4 m² to 9.1 × 10⁻¹ 4 The porosity is comparable to that of conventional porous permeable steel, with a tensile strength ranging from 760 MPa to 1150 MPa and a maximum DC relative permeability of 135 to 410. Under constant excitation current at 100 kHz, the surface induction heating rate of the exhaust insert is 7.0 °C / s to 11.9 °C / s, while that of conventional porous permeable steel with the same porosity is 450 MPa, 28 °C / s, and 3.2 °C / s, respectively. In a further preferred range with additive dosages of 8% to 12%, the tensile strength is not less than 985 MPa, the maximum DC relative permeability is not less than 320, and the surface induction heating rate is not less than 10.8 °C / s. The original mutual constraints between porosity and mechanical and electromagnetic properties are thus alleviated.

[0020] Because the two-dimensional layered ceramic nanosheets do not wet the transient liquid phase, they are repelled to the three-phase interface and adhere to the inner wall of the pores during melt aggregation and solidification. The inner surface of the micron-scale interconnected channels completes low surface energy modification in a single sintering process, eliminating the need for secondary post-treatments such as chemical vapor deposition, physical vapor deposition, or impregnation passivation. After 30,000 cycles of polyetherimide melt injection molding, the material's air permeability retention rate is 74.8% to 98.0%, with a nanosheet mass fraction of no less than 95.2% when it is 2% to 5%, compared to 25.4% for conventional porous permeable steel. According to the neutral salt spray test, the time for red rust to appear on the surface is extended from 48 hours for conventional porous permeable steel to over 96 hours. When the nanosheet mass fraction is 2% to 5% and the additive dosage is 8% to 12%, no red rust appears even after 720 hours of testing.

[0021] In addition, nickel and copper in the transient liquid phase diffuse into the deep layers of the matrix lattice along the sintering neck during the second holding stage. While forming a solid solution strengthening bridging layer through isothermal solidification, they reduce the critical cooling rate of the matrix. The porous body can obtain a uniform martensitic structure along the cross section without the use of quenching oil and only under high-pressure nitrogen quenching. The apparent hardness difference between the surface and the core is reduced from 4.2 HRC of conventional porous permeable steel to less than 2 HRC. The macroscopic hardness is not reduced due to the introduction of soft magnetic components. Attached Figure Description

[0022] Figure 1 This is a schematic diagram of the structure of a single composite particle of the multi-component composite alloy additive in a preferred embodiment of the present invention; Figure 2 This is a schematic diagram of the microstructure of the sintered neck and pore inner wall of the porous and breathable plastic mold steel in a preferred embodiment of the present invention. Figure 3 This is a temperature regime curve for stepped heating sintering in a preferred embodiment of the present invention.

[0023] Figure labeling: 1. Soft magnetic alloy nanoparticles; 2. Transient liquid phase alloy shell; 3. Two-dimensional layered ceramic nanosheets; 4. Plastic mold steel matrix powder particles; 5. Sintering neck; 6. Connecting pores; 7. Low surface energy isolation layer; 8. Solid solution reinforced diffusion region; 9. Cobalt-enriched soft magnetic phase region. Detailed Implementation

[0024] like Figure 1 As shown, the multi-component composite alloy additive of this invention is not a simple combination of several powders, but a composite pre-constructed at the single-particle scale. It has a soft magnetic alloy nanoparticle 1 as its core, with a continuous transient liquid phase alloy shell 2 deposited by chemical plating on the outer surface. The outermost surface of the shell is coated with two-dimensional layered ceramic nanosheets 3. By incorporating this additive into a plastic mold steel matrix powder in a specific ratio, followed by cold isostatic pressing and stepped temperature sintering, liquid phase strengthening of the sintering neck, targeted enrichment of the soft magnetic phase, and low surface energy modification of the pore walls can be simultaneously achieved within the porous framework. The resulting structure is as follows: Figure 2 As shown, the powder particles 4 of the plastic mold steel matrix are interconnected by widened sintering necks 5, the gaps between the particles are retained as interconnected pores 6, the inner walls of the pores are covered with a low surface energy isolation layer 7, the two sides of the sintering neck are solid solution strengthened diffusion zones 8, and cobalt-enriched soft magnetic phase zones 9 are distributed around the direct contact area remaining between adjacent particles in the sintering neck. The following description, combined with specific raw materials, processes and test data, further illustrates the present invention. The embodiments described are only for explaining the present invention and do not constitute a limitation on the scope of protection of the present invention.

[0025] The matrix powder, SUS420, is a martensitic stainless steel powder prepared by argon atomization after vacuum induction melting. It was provided by AVIC MITEK Powder Metallurgy Technology (Beijing) Co., Ltd., with a particle size range of 15 μm to 45 μm, a median diameter of 30 μm, a sphericity of not less than 0.9, and an oxygen content of not more than 0.03% by mass. Its solidus temperature, determined by differential thermal analysis, is 1385℃. Some embodiments used H13 hot work die steel powder of the same particle size, provided by Hunan Yingjie High-Tech Co., Ltd. The soft magnetic alloy nanoparticles are Fe50Co48V2 nanoparticles, prepared from a master alloy of the corresponding composition via hydrogen plasma metal reaction, with a particle size of 50 nm to 200 nm and a purity of not less than 99.5%. The two-dimensional layered ceramic nanosheets are hexagonal boron nitride nanosheets, prepared by liquid-phase exfoliation of bulk hexagonal boron nitride with N-methylpyrrolidone followed by centrifugal fractionation. The thickness is no greater than 5 nm, the transverse sheet diameter is 200 nm to 500 nm, and the purity is no less than 99%. The molding binder is polyvinyl butyral, prepared as a 5% (w / w) anhydrous ethanol solution. An anhydrous system was chosen to avoid oxidation of the phosphorus-containing nickel-based shell and the nano-iron-cobalt phase in water and hot air during the granulation process, thus aligning with the surface cleanliness requirements of the dry mixing and sintering processes.

[0026] The composite additive was synthesized as follows: Fe50Co48V2 nanoparticles were successively treated with stannous chloride sensitizing solution and palladium chloride activating solution, and washed with deionized water until neutral. They were then dispersed in a chemical plating bath under ultrasonic and mechanical stirring. The plating bath used nickel sulfate and copper sulfate as metal sources, sodium hypophosphite as a reducing agent, and contained lactic acid and propionic acid as complexing and buffering components. The pH was adjusted to 5.0 with ammonia, and the bath temperature was maintained at 85℃. The shell layer was deposited to the set mass fraction by controlling the ratio of the total metal ions in the plating bath to the amount of nanoparticles. After plating, the particles were centrifuged, washed, and dried in a vacuum at 60℃ to obtain coated particles. Inductively coupled plasma atomic emission spectrometry (ICP-AES) analysis showed that the shell layer composition was P 10.2%, Cu 9.8%, and the balance Ni. Differential scanning calorimetry (DSC) under argon protection and a heating rate of 10℃ / min determined that the solidus temperature and liquidus temperature of the shell alloy after crystallization were 866℃ and 895℃, respectively. Using the seat drop method at a vacuum degree of 5×10⁻ 4The wetting angle of the alloy melt on a hot-pressed hexagonal boron nitride dense substrate was measured at Pa and 950℃. The equilibrium contact angle was 138°, which is typical of a non-wetting system. Subsequently, the coated particles and hexagonal boron nitride nanosheets were loaded into a nylon can at a set ratio, and zirconia grinding balls were added. After argon purging three times, the mixture was dry-mixed in a planetary mixer at a revolution speed of 200 r / min. The net running time was 45 min, and the mixer was stopped for 5 min every 15 min to suppress the temperature rise of the can. Scanning electron microscopy revealed that the coated particles of the composite powder were nearly spherical. Transmission electron microscopy showed that continuous shells with a thickness of 30 nm to 60 nm could be distinguished, with uniform contrast and equal thickness distribution along the core outline. The nanosheets were attached to the surface of the coated particles in an adherent manner. Laser particle size analysis showed that the median diameter of the coated particles was 0.24 μm, which was consistent with the geometric relationship between the core particle size and the shell thickness. The composite powder existed as soft aggregates with a median diameter of about 1.5 μm, which could be deagglomerated during subsequent wet mixing.

[0027] The composite additive was ultrasonically dispersed in anhydrous ethanol for 30 min and then mixed with the matrix powder. Mixing continued for 4 h in a gravity-free three-dimensional mixer. All samples were mixed using the same equipment and mixing regime. After mixing, the powder was dried in a vacuum at 40℃, then sprayed with a polyvinyl butyral ethanol solution for granulation. The binder solid content was controlled at 0.8% of the total mass of the mixed powder. The powder was then vacuum dried at 40℃ for 4 h under argon protection and sieved through a 250 μm sieve for granulation. The granulated powder was packed into a rubber sleeve, vacuum sealed, and then cold isostatically pressed. After holding the pressure for 5 min, the pressure was released and the sample was demolded to obtain a green body of the specified shape. The green body was placed in a three-zone temperature-controlled vacuum atmosphere furnace with a molybdenum heating element. The temperature uniformity of the furnace's uniform temperature zone was no higher than ±5℃. Thermocouples were pre-embedded inside the sample to verify that the temperature difference between the inside and outside of the sample converged to within 3℃ during the holding period. Figure 3 The temperature regime shown is as follows: In a flowing reducing atmosphere containing 5% hydrogen and the balance argon, the temperature is increased to 450℃ at a rate of 1℃ / min and held for 3 hours to complete degreasing; a vacuum is then applied to a pressure better than 5×10⁻³ Pa, followed by heating to the first holding stage temperature at a rate of 3℃ / min and holding; high-purity argon is then refilled to a partial pressure of 500 Pa, and the temperature is increased to the second holding stage temperature at a rate of 5℃ / min and held; subsequently, the furnace is cooled to room temperature. Carbon-sulfur analysis shows that the residual carbon content of the degreased green body is less than 0.03% by mass.

[0028] Under this temperature regime, the evolution within the material exhibits a clear temporal sequence. At the end of degreasing, the spherical matrix powder particles 4 retain the original geometric gaps shaped by low-pressure isostatic pressing, forming the topological prototype of interconnected pores 6. Meanwhile, the soft aggregates of the composite additives preferentially adhere to the surface of the matrix particles and accumulate near the particle contact area during mixing and granulation. When the temperature exceeds the liquidus line of the shell alloy, the shell 2 melts, and the previously coated Fe-Co-V nanoparticles immediately transform into a dispersed phase suspended in the melt. Based on the density of each phase, the volume fraction of suspended particles in the melt at this point is approximately 20%, and according to the Krieger-Dougherty relationship, its relative apparent viscosity is approximately 1.8 times the intrinsic viscosity of the melt, indicating that the melt still possesses good flowability. The melt forms a saddle-shaped liquid surface at the narrow neck where the matrix particles come into contact. According to the Laplace equation, the principal radius of curvature at this point is extremely small, resulting in a significant local negative pressure inside the melt, causing the melt to spontaneously converge towards the neck. Suspended particles migrate with the flow field under the influence of viscous drag and accumulate densely at the neck. At the same time, the wetting contact angle between the nickel-based melt and hexagonal boron nitride is much greater than 90°. In order to minimize the total interfacial free energy, the system continuously pushes the nanosheets 3 to the three-phase interface of solid, liquid and pores during the melt contraction and convergence process, keeping them outside the melt front. After heating to the second holding stage, phosphorus in the melt rapidly diffuses into the matrix lattice and grain boundaries along the concentration gradient. Accompanying this is iron in the matrix dissolving into the bonding region in the opposite direction. The asymmetry of diffusion flux triggers the Kirkendall effect. Due to the continuous loss of phosphorus and the continuous dissolution of iron, the composition of the bonding region shifts towards the Ni-Cu-Fe solid solution side. This solid solution is mainly composed of nickel with a melting point of 1455℃. Its solidus temperature rises rapidly with the shift in composition. When the holding temperature exceeds 1150℃, the bonding region changes from liquid to solid under isothermal conditions, forming a solid solution-enhanced diffusion region 8. It should be noted that during this heat preservation process, Fe-Co-V particles with a size in the nanometer range undergo partial dissolution and Ostwald ripening, ultimately forming a cobalt-enriched Fe-Co-Ni solid solution region with a particle size of 0.3 μm to 1.5 μm at the sintering neck 5, with some incompletely dissolved fine particles remaining, constituting the cobalt-enriched soft magnetic phase region 9. Since the melt can only enter the wedge-shaped gaps around the direct contact area of ​​adjacent particles, the cobalt-enriched soft magnetic phase region 9 is distributed in a ring-like pattern around this direct contact area within the solidified bridging layer. Its cobalt content is still significantly higher than that of the matrix, and its saturation magnetization is also higher than that of the surrounding martensitic phase. The magnetic flux pathway in the neck is not lost due to the coarsening of the particles, and the trace amount of vanadium in it delays this coarsening process by enriching at the interface. During the cooling and shrinkage stage, the nanosheets previously pushed to the melt front are stably attached to and fixed on the inner surface of the solidified pores, forming a low surface energy isolation layer 7.

[0029] It should be noted that the temperature of the second holding stage is significantly lower than that of conventional powder metallurgy sintering. This invention does not rely on the solid-state sintering of the matrix powder itself to establish the skeleton's connection, and the holding temperature of 1150℃ to 1170℃ is not intended to densify the matrix. The purpose of this temperature range is to provide the necessary diffusion kinetics for isothermal solidification. The connection strength of the skeleton is mainly borne by the bridging layer after isothermal solidification. Therefore, it is unnecessary to raise the temperature to the conventional range of 1250℃ to 1350℃. This is precisely why this invention can achieve a strength far exceeding that of conventional porous permeable steel while maintaining a porosity of approximately 20%. Metallographic observation shows that the neck width of the sample of the present invention can reach 0.35 to 0.45 times the diameter of the matrix particles, of which about three-quarters of the width is contributed by the solidified bridging layer, and the remaining about one-quarter is the direct contact area between adjacent particles. In contrast, the neck width of the comparative sample sintered at 1250°C for 3 h using conventional processes is only about 0.12 times the diameter of the particles. On the tensile fracture surface, the crack path of the sample of the present invention mainly penetrates the matrix particles and exhibits a mixture of dimples and cleavage. No continuous cracking along the interface of the bridging layer or intergranular fracture characteristics were observed.

[0030] The phase stability of the pore wall modification layer was specifically verified after sintering. The characteristic diffraction peaks of hexagonal boron nitride were still clearly detected in the X-ray diffraction spectrum, and their relative intensity showed no significant decrease compared to before sintering. Transmission electron microscopy combined with electron energy loss spectroscopy showed that the nanosheets on the pore walls maintained a complete layered structure, with a transition layer less than 5 nm thick on the side in contact with the substrate. Chromium, boron, and nitrogen signals were detected at this point, indicating that a limited reaction occurred at the interface, generating a small amount of chromium borides and nitrides. The presence of these reaction products actually anchored the nanosheets firmly to the pore walls, preventing them from easily detaching under high-pressure gas flow. However, the reaction depth was limited by solid-state diffusion, preventing overall decomposition of the nanosheets. Based on this interfacial reaction, the upper limit of the temperature in the second holding stage was set at 1170°C to avoid the eutectic temperature of 1174°C for iron and borides, thus preventing grain boundary liquefaction caused by accidental release of boron.

[0031] The performance tests were conducted according to the following methods. The total porosity was calculated from the measured skeleton density and the theoretical density. The theoretical density was obtained by summing the volumes of the densities of each phase, where SUS420 was 7.75 g / cm³, H13 was 7.80 g / cm³, Ni-Cu-P shell was 8.00 g / cm³, Fe-Co-V was 8.10 g / cm³, and hexagonal boron nitride was 2.27 g / cm³. Taking Example 1 as an example, 100 g of matrix powder was mixed with 10 g of additives, including 7.8 g of shell, 2.0 g of soft magnetic phase, and 0.2 g of nanosheets, with a total volume of 14.21 cm³. The theoretical density was 7.74 g / cm³, and the measured skeleton density was 6.15 g / cm³. Based on this, the total porosity was 20.5%. The open porosity was determined according to GB / T 5163-2006 using the Archimedes method, and the open porosity of each sample accounted for more than 95% of its total porosity. Gas permeability was measured using a Darcy permeameter. The sample was a circular disc with a diameter of 25 mm and a thickness of 5.0 mm. The medium was dry nitrogen at room temperature, and the outlet pressure was atmospheric pressure. The inlet gauge pressure was taken at four points: 0.02 MPa, 0.03 MPa, 0.05 MPa, and 0.08 MPa. The apparent permeability at each point was calculated using the Darcy formula for compressible fluids. Then, a linear regression was performed on the apparent permeability against the reciprocal of the average pressure, and extrapolation was performed until the reciprocal of the average pressure approached zero. After deducting the Klinkenberg slip effect, the intrinsic permeability was obtained. The calculated pore Reynolds number at each measuring point was less than 1, indicating that the flow was in the Darcy linear region, and the inertial term was negligible. The heat treatment process was as follows: SUS420 matrix samples were vacuum heated at 1030℃ for 30 min, then quenched under high pressure with 0.6 MPa high-purity nitrogen, and then vacuum tempered twice at 200℃ for 2 h each time; H13 matrix samples were austenitized and quenched at 1030℃, then vacuum tempered twice at 560℃ for 2 h each time. All samples were quenched under high pressure instead of oil quenching because this material is a highly interconnected porous body with an open porosity of about 20%. Quenching oil would fill the pores under capillary action and could not be completely removed. Residual oil would not only distort the permeability and injection molding aging data, but would also crack and coke during high-temperature injection molding, artificially causing pore blockage. Tensile strength specimens were directly formed by pressing and sintering according to GB / T 7963-2005, without surface machining, to avoid surface sealing and coating caused by machining. Testing was conducted at room temperature according to GB / T 228.1-2021, with a loading rate of 2 mm / min. The arithmetic mean of 5 specimens per group was taken. Apparent hardness was determined according to the Rockwell hardness C scale of GB / T 9097-2016. The reading at 0.5 mm from the surface was taken as the reported value, and a separate reading was taken at the core to examine the uniformity of cross-sectional hardness. The arithmetic mean of 10 points at each location was taken.The magnetic properties were determined according to GB / T13012-2008 using a ring-shaped sample with an outer diameter of 25 mm, an inner diameter of 15 mm, and a height of 5 mm, and a double winding method, under DC magnetization conditions. The resistivity was determined by the four-probe method, and in Example 1 it was 1.28 × 10⁻. 6 Ω·m, Comparative Example 1 is 1.35 × 10⁻ 6 Based on the Ω·m, the skin depths of the two at 100 kHz were calculated to be approximately 0.10 mm and 0.35 mm, respectively. Induction heating performance was assessed using actual inserts. The sample was machined into a 40 mm × 40 mm × 20 mm insert and placed 2 mm below a planar induction coil operating at 100 kHz with a constant excitation current of 320 A. The average heating rate of the surface from 50°C to 150°C was recorded using an infrared thermometer. The thermal diffusivity was measured at 25°C using the laser scintillation method according to GB / T 22588-2008, and converted to thermal conductivity using specific heat capacity and density. Long-term service anti-clogging performance was assessed through simulated injection molding aging tests. The sample was machined into a venting insert and placed in a mold. Using polyetherimide as the raw material, 30,000 injection cycles were performed under the same conditions: barrel temperature 380°C, mold temperature 160°C, and injection pressure 120 MPa. Gas flow rate was measured before and after the test, and the air permeability retention rate was calculated. The corrosion resistance was tested according to GB / T 10125-2021 using a neutral salt spray test. The sodium chloride solution concentration was 50 g / L ± 5 g / L, the pH value of the collected solution was controlled between 6.5 and 7.2, the temperature inside the chamber was 35℃, and the time when obvious red rust appeared on the sample surface was recorded. The test was conducted uniformly for 720 h. Example 1

[0032] A composite additive was synthesized based on a mass ratio of 20% soft magnetic alloy nanoparticles, 2% two-dimensional layered ceramic nanosheets, and the remainder being a transient liquid-phase alloy shell. 100 parts of SUS420 matrix powder were mixed with 10% of this additive by mass, along with 0.8% polyvinyl butyral, and granulated under cold isostatic pressing at 200 MPa. The first holding stage during sintering was set at 940℃ for 2 h, and the second holding stage was set at 1150℃ for 3 h. The mass fraction of phosphorus in the sintered body was determined to be 0.70% by inductively coupled plasma atomic emission spectrometry. Example 2

[0033] The composition and synthesis process of the composite additive were the same as in Example 1, but its dosage was reduced to 5% of the matrix powder mass. The remaining molding and sintering conditions were the same as in Example 1. The mass fraction of phosphorus in the sintered body was 0.37%. Example 3

[0034] The composition and synthesis process of the composite additive are the same as in Example 1, but its dosage is increased to 15% of the matrix powder mass. The remaining molding and sintering conditions are the same as in Example 1. The mass fraction of phosphorus in the sintered body is 1.02%. Example 4

[0035] A composite additive was synthesized based on the following mass ratio: 25% soft magnetic alloy nanoparticles, 5% two-dimensional layered ceramic nanosheets, and the balance being a transient liquid-phase alloy shell. The additive was used at 12% of the matrix powder mass. The cold isostatic pressing pressure was increased to 250 MPa. The first holding stage was at 950℃ for 2.5 h, and the second holding stage was at 1165℃ for 3 h. All other conditions were the same as in Example 1. The phosphorus mass fraction in the sintered body was 0.75%. Example 5

[0036] A composite additive was synthesized according to the mass ratio of 12% soft magnetic alloy nanoparticles, 3% two-dimensional layered ceramic nanosheets, and the balance being a transient liquid-phase alloy shell. The additive was used at 8% of the mass of the matrix powder. The matrix powder was replaced with H13 hot-work die steel powder of the same particle size. The cold isostatic pressing pressure was reduced to 150 MPa. The first holding stage was at 940℃ for 2 h, and the second holding stage was at 1160℃ for 2.5 h. The remaining conditions were the same as in Example 1. The phosphorus mass fraction in the sintered body was 0.63%. Example 6

[0037] A composite additive was synthesized according to the mass ratio of 20% soft magnetic alloy nanoparticles, 1% two-dimensional layered ceramic nanosheets, and the balance being a transient liquid phase alloy shell. Specifically, the interface modification phase was taken up to the lower limit of the range defined in the claims, and its dosage was 10% of the matrix powder mass. The remaining molding and sintering conditions were completely consistent with those of Example 1. The mass fraction of phosphorus in the sintered body was 0.71%. Contrast 1

[0038] Without using the composite additives of this invention, urea microspheres equivalent to 10% of the mass of the SUS420 matrix powder were incorporated as a pore-forming agent according to industry practice, and cold isostatic pressing was performed at 200 MPa. To eliminate interference from differences in the debinding process, this comparative example adopted the same slow debinding process as the example, i.e., heating to 450°C at 1°C / min and holding for 3 h in a hydrogen-containing reducing atmosphere, followed by heating to the conventional sintering temperature of martensitic stainless steel powder, 1250°C, at 5°C / min and holding for 3 h, relying on the pyrolysis and escape of urea in the low-temperature region to form pores. Contrast 2

[0039] The types, chemical compositions, and relative proportions of all components in the system were exactly the same as in Example 1. The only difference was that the soft magnetic alloy nanoparticles were not chemically coated. Instead, Ni-Cu-P alloy powder, Fe-Co-V nanoparticles, and hexagonal boron nitride nanosheets were directly mixed with the matrix powder using the same three-dimensional mixer and the same mixing regime. The molding and sintering regimes were completely consistent with those in Example 1. Contrast 3

[0040] The synthesis process of the composite additive is the same as that in Example 1, but the hexagonal boron nitride nanosheets are removed during dry mixing. The additive retains only the binary structure of soft magnetic alloy nanoparticles and transient liquid phase alloy shell. The core-shell mass ratio is prepared at 20:80 to keep the total amount unchanged. The amount used is still 10% of the mass of the matrix powder. The molding and sintering process is completely consistent with that in Example 1. Contrast 4

[0041] The composition, synthesis process and dosage of the composite additive are the same as in Example 1. The only difference is that the sintering process adopts the conventional practice of existing technology, the first heat preservation stage is cancelled, and after degreasing, the temperature is continuously increased at a rate of 10℃ / min to directly reach 1150℃ and held for 3 hours. Contrast 5

[0042] The types, particle sizes, and mass ratios of the soft magnetic alloy nanoparticles and hexagonal boron nitride nanosheets in the composite additive were the same as in Example 1. The only difference was that the shell layer was replaced with a Cu-Mn-Ni alloy with copper as the main metal, comprising 60% Cu, 32% Mn, and 8% Ni by mass percentage. This alloy used manganese as a melting point reducing element. Its solidus temperature was 866°C and liquidus temperature was 891°C, determined by differential scanning calorimetry. These temperatures were basically consistent with the melting range of the shell alloy used in Example 1. However, the melting point of copper, the metal element with the highest mass fraction, was only 1085°C, which was lower than the temperature of the second holding stage. The forming and sintering regimes were completely consistent with those of Example 1, namely, holding at 940°C for 2 hours in the first holding stage and holding at 1150°C for 3 hours in the second holding stage.

[0043] Table 1. Composition and main process parameters of each embodiment and comparative example.

[0044] Table 2 Performance test results of each embodiment and comparative example

[0045] Comparing Example 1 with Comparative Example 1, which represents conventional practices, the total porosity of Example 1 was 20.5% and 22.5%, respectively, and the gas permeability was 6.2 × 10⁻¹, respectively.4 m² and 8.8 × 10⁻¹ 4 The two samples, with a diameter of m², are on the same order of magnitude and comparable in terms of basic exhaust function. However, the tensile strength of Example 1 increased from 450 MPa to 985 MPa, the maximum relative permeability under DC current increased from 28 to 320, and the surface induction heating rate of the actual insert increased from 3.2℃ / s to 10.8℃ / s. The microscopic reason for this result lies in the completely different morphology and composition of the sintered necks of the two samples. The sintered neck of Comparative Example 1 is formed only by solid-phase diffusion at the particle point contact, and its width is only about 0.12 times the particle diameter. Residual oxide films from the original powder surface are also retained on both sides. Under load, the stress is highly concentrated in this section, and during magnetization, the magnetic flux is also forced to be transmitted entirely through this section and saturates first. Therefore, the effective permeability of the porous body is dominated by the proportion of the neck section. According to the relationship between the effective medium of randomly distributed non-magnetic inclusions, the reduction of the effective permeability by about 20% of the pores is only about one-quarter, which is far from enough to explain the value of 28. It can be seen that the magnetic neck contraction formed by the neck section is the dominant factor for the low permeability, which is also the focus of this invention. In Example 1, the melt converges towards the neck under Laplace negative pressure, widening the point contact filling to 0.35 to 0.45 times the particle diameter. Geometrically, the stress concentration factor and magnetoresistance decrease simultaneously. The diffusion of phosphorus and nickel not only dissolves and breaks down the original oxide film but also forms solid solution-strengthened transition zones on both sides of the neck, changing the fracture path from along the neck interface to through the matrix. The cobalt-rich phase transported to the neck by the melt has a higher saturation magnetization than the martensitic matrix, further reducing the magnetoresistance of the neck and delaying its saturation. The porosity remains almost unchanged while the strength and permeability increase simultaneously, indicating that the performance improvement does not come from densification but from the reconstruction of the spatial distribution of the material within the porous framework. Regarding induction heating, the skin depth decreases according to the reciprocal of its square root after the increase in permeability. Based on measured resistivity, the skin depth of Example 1 at 100 kHz is approximately 0.10 mm, while that of Comparative Example 1 is approximately 0.35 mm. The energy deposition in the former is concentrated in a thinner surface layer, which explains the difference in heating rate. The change in thermal conductivity also supports the above judgment: according to the Maxwell-Eucken relationship, with a dense matrix of 24.9 W / (m·K), the theoretical upper limit corresponding to 20.5% porosity is about 18.0 W / (m·K). The measured value of 16.8 W / (m·K) in Example 1 is close to this upper limit, indicating that the shrinkage thermal resistance of the neck has been basically eliminated. However, the theoretical upper limit of Comparative Example 1 at 22.5% porosity is about 17.3 W / (m·K), while its measured value is only 8.6 W / (m·K), less than half of the upper limit, which is due to the shrinkage thermal resistance caused by the narrow neck.

[0046] Examples 2 and 3 illustrate the boundary effects of additive dosage. In Example 2, the dosage was 5%, resulting in insufficient total melt volume, which only formed encapsulation at some contact points, insufficient neck widening, and a drop in tensile strength to 760 MPa. More sensitively, the magnetic properties were affected; the enrichment of the transported soft magnetic phase at the neck was insufficient to significantly reduce magnetoresistance, with a maximum relative permeability of only 135. Simultaneously, the absolute amount of nanosheets was only 0.1% of the matrix powder mass, leading to incomplete pore wall coverage and a decrease in permeability retention to 78.5%. In Example 3, the dosage was 15%, with excess melt overflowing into the bulk air gap after filling the contact points, reducing the total porosity to 13.2% and the gas permeability to 1.4 × 10⁻¹. 4 Although the strength and permeability of the sample increased further to 1150 MPa and 410 m², respectively, the venting capacity was significantly weakened. Simultaneously, the phosphorus mass fraction in the sintered body increased to 1.02%, exceeding the conventional control level for this element in powder metallurgy iron-based materials. The measured elongation after fracture decreased by approximately one-third compared to Example 1, and intergranular characteristics appeared locally on the fracture surface. These two boundary effects together constitute the basis for further optimizing the dosage to 8% to 12% and limiting the mass fraction of the demelting element in the sintered body to 0.3% to 0.8%. Examples 4 and 5 illustrate that different performance balances can be obtained within this range by synergistically adjusting the proportions and temperature regime. In Example 4, the soft magnetic phase and nanosheets were simultaneously taken to the upper limit, and the second heat treatment stage was raised to 1165℃, resulting in high strength and permeability while the permeability decreased moderately. In Example 5, the soft magnetic phase was taken to the lower limit, and an H13 matrix was used instead. Tempering was performed at 560℃ according to its secondary hardening zone, maintaining a porosity of 21.9% and a permeability of 8.0×10⁻¹. 4 While achieving a permeability of m², it still obtains a strength of 880 MPa and a maximum relative permeability of 268, indicating that the mechanism of action of this invention is not limited to a single matrix grade.

[0047] Comparative Example 2 best illustrates the necessity of pre-constructing a coating structure. Its macroscopic chemical composition is completely consistent with Example 1, and the mixing equipment and mixing regime are also the same as in all examples. The only difference is that the soft magnetic alloy nanoparticles are not chemically coated. Due to the lack of spatial constraints established by the shell, the Fe-Co-V nanoparticles and Ni-Cu-P alloy powder are distributed in isolation in the green blank. After heating, although the melt can still converge towards the neck due to capillary action, it can only carry a small number of soft magnetic particles that it encounters by chance on the flow path. Most of the soft magnetic nanoparticles either remain near the pore walls far from the neck, or, due to the lack of the protective shell, their surfaces oxidize during heating and interdiffusion with the iron atoms in the matrix, losing their compositional advantages. Ultimately, the maximum relative permeability is only 62. At the same time, the free hexagonal boron nitride nanosheets directly remain between the matrix particles where the melt has not yet flowed, which instead blocks some of the contact points that should have formed a metallurgical bond, resulting in a tensile strength of only 580 MPa. This indicates that the effect of the present invention does not come from the macroscopic superposition of components, but from the coupling of the arrangement of components in the microscopic space with the sintering sequence.

[0048] Comparative Example 3, along with Examples 6 and 1, constitutes a complete gradient regarding the interface-modified phase, containing a potentially overlooked piece of reverse evidence. Comparative Example 3, which removed hexagonal boron nitride, exhibited porosity, strength, and permeability comparable to Example 1, indicating that this component contributes almost nothing to mechanical and electromagnetic properties. However, after 30,000 cycles of polyetherimide injection molding, its permeability retention was only 32.7%. This is because the pore walls remain exposed metal free surfaces, where unpaired dangling bonds between iron and chromium atoms constitute high-energy chemisorption sites. Phenolic and free radical fragments generated by the pyrolysis of polyetherimide at 380°C readily undergo chemisorption and further condensation and carbonization at these sites. The deposited layer thickens layer by layer over tens of thousands of cycles, causing rapid scaling of the originally ten-micrometer-wide pores. More noteworthy is the salt spray test result of Comparative Example 3, where red rust appeared in 36 hours, shorter than the 48 hours of Comparative Example 1 without any alloying additives. This reverse change is consistent with the compositional evolution after the introduction of additives: after adding 10% additives, the nominal chromium content in the matrix is ​​diluted from 13% to approximately 11.8%, while approximately 0.70% phosphorus is introduced. Phosphorus weakens the integrity of the passivation film, and the nickel-rich and copper-rich bridging layer and the iron-chromium matrix also form an electrical couple in the chloride medium, promoting selective corrosion at the neck. Therefore, the improved corrosion resistance reported in this invention does not come from the introduction of alloying elements, but entirely from the in-situ formation of a hydrophobic isolation layer on the pore walls, which blocks the capillary penetration of the chloride-containing liquid film. Removing this isolation layer not only fails to improve corrosion resistance, but actually results in lower corrosion resistance than the unmodified porous body. In Example 6, the nanosheet content was set to the lower limit of 1%, resulting in insufficient coverage to form a continuous barrier. The retention rate was 74.8%, and the red rust time was 96 h, falling between Comparative Example 3 and Example 1. In Example 1, the nanosheet content was 2%, and the pore walls were covered by a basically continuous layered structure, achieving a retention rate of 96.5%, with no red rust appearing even after 720 h of testing. This gradient relationship directly confirms the existence of the mechanism by which the non-wetting phase is repelled to the three-phase interface and forms a film in situ during the solidification of the melt, and also confirms that this isolation layer is the only source of improved corrosion resistance.

[0049] Comparative Example 4 further illustrates the decisive role of temperature regime in the above mechanism. Its composition is exactly the same as that of Example 1, except that the first holding stage is omitted and the temperature is rapidly and continuously increased according to conventional methods. At a heating rate of 10°C / min, the shell melts and is in a flowable state for only a few minutes before entering the high-temperature zone. The melt has not yet completed capillary convergence towards the neck before being absorbed by the matrix due to rapid interdiffusion of elements. Isothermal solidification occurs prematurely in situ, and the soft magnetic phase is thus largely retained in its original position, with a maximum relative permeability of only 96. The hexagonal boron nitride nanosheets are also engulfed and encapsulated in the solidified structure because the melt does not have time to be pushed out, and fail to migrate to the pore walls, resulting in a permeability retention rate of only 61.2%. It can be seen that the kinetic time window provided by the first holding stage is a prerequisite for the full occurrence of melt targeted transport and interface self-assembly.

[0050] Comparative Example 5 thermodynamically verifies that limiting the melting point of the main metal in the shell alloy is not superfluous. The Cu-Mn-Ni alloy used in this comparative example and the Ni-Cu-P alloy used in Example 1 both have a solidus temperature of 866°C and liquidus temperatures of 891°C and 895°C respectively, with almost overlapping melting ranges. Therefore, they can form a control system under the same sintering regime, changing only the main metal. Melt formation, capillary aggregation, and the transport and displacement of the soft magnetic and non-wetting phases still occur in the first holding stage. Therefore, its maximum relative permeability still reaches 235, and the permeability retention rate remains at 88.2%, indicating that the aforementioned transport mechanism itself is unrelated to the choice of the main metal. The real difference occurs in the second holding stage: copper has a melting point of 1085°C. At a holding temperature below 1150°C, even if all manganese and nickel diffuse away, the composition of the bonding zone can only approach that of a copper-based solid solution containing a small amount of iron. Its melting point remains below the ambient temperature, making isothermal solidification thermodynamically impossible. The consequences are threefold. First, the bonding zone remains liquid throughout the high-temperature insulation period, lacking the support of a solid-phase framework. Under the influence of surface tension, the liquid film continues to spread into the air gap, causing the total porosity to drop to 16.8% and the permeability to drop to 2.9 × 10⁻¹. 4 Secondly, liquid copper penetrates along the austenite grain boundaries of the matrix and forms a continuous low-melting-point grain boundary film after cooling, inducing typical liquid metal embrittlement, with a tensile strength of only 520 MPa and a fracture surface exhibiting obvious intergranular characteristics; thirdly, the residual copper-rich phase is a soft phase and does not undergo martensitic transformation during quenching, causing the apparent hardness to drop to 46.8 HRC, the lowest value among all samples. Therefore, the melting point of the main metal must be higher than the temperature of the second holding stage for the isothermal solidification and solid solution strengthening bridging layer to truly form, which is also the basis for the limitation of this feature in claim 1.

[0051] From the cross-sectional hardness distribution, the apparent hardness of the sample in Example 1 at 0.5 mm from the surface and at the core were 50.5 HRC and 49.2 HRC, respectively, with a difference of 1.3 HRC. In contrast, the hardness at the corresponding positions in Comparative Example 1 were 47.5 HRC and 43.3 HRC, respectively, with a difference of 4.2 HRC. This difference stems from the fact that nickel and copper in the shell diffused into the deeper layers of the matrix lattice along the neck during the second heat treatment stage. As elements that reduce the critical cooling rate, they delayed the transformation of pearlite and bainite, allowing the porous material to achieve sufficient martensitic transformation even under conditions of only 0.6 MPa nitrogen quenching and a significantly lower cooling rate in the core than in the surface layer, thus improving hardenability. The soft magnetic phase was confined to the sintered neck, a region with a very small volume, and was encapsulated by the solidified solid solution. Its dilution effect on macroscopic hardness was negligible, and the material did not experience a decrease in hardness due to the introduction of the soft magnetic component. As can be seen from the data in Table 2, this invention achieves a balance between porosity, strength, magnetic permeability and anti-clogging ability by pre-arranging the components at the single particle scale and using temperature regime to make each component play its role sequentially on the time axis. This effect cannot be obtained by directly mixing equal amounts of components, conventional sintering regime, or low-melting-point brazing filler metal with insufficient melting point of the main metal.

[0052] The above description is merely a preferred embodiment of the present invention and is not intended to limit the invention. Various modifications and variations can be made to the present invention by those skilled in the art. Any modifications, equivalent substitutions, improvements, etc., made within the spirit and principles of the present invention should be included within the scope of protection of the present invention.

Claims

1. A multi-component composite alloy additive for high hardenability plastic mold steel, characterized in that, The multi-component composite alloy additive is a composite powder. Each particle of the composite powder comprises soft magnetic alloy nanoparticles, a transient liquid phase alloy shell coating the outer surface of the soft magnetic alloy nanoparticles, and two-dimensional layered ceramic nanosheets attached to the outer surface of the transient liquid phase alloy shell. The transient liquid phase alloy contains at least one melting point reducing element selected from phosphorus, boron, and silicon. The liquidus temperature of the transient liquid phase alloy is 850°C to 920°C, and the melting point of the metal element with the highest mass fraction in the transient liquid phase alloy is not lower than 1400°C. The wetting contact angle of the melt of the transient liquid phase alloy on a dense substrate of the same material as the two-dimensional layered ceramic nanosheets is greater than 90°. The wetting contact angle is a value measured by the seat drop method under the conditions of temperature 950°C and vacuum degree not higher than 10⁻³ Pa.

2. The multi-component composite alloy additive according to claim 1, characterized in that, Based on the total mass of the multi-component composite alloy additives, the mass fraction of the soft magnetic alloy nanoparticles is 12% to 25%, the mass fraction of the two-dimensional layered ceramic nanosheets is 1% to 5%, and the balance is the transient liquid phase alloy and unavoidable impurities.

3. The multi-component composite alloy additive according to claim 1 or 2, characterized in that, The soft magnetic alloy nanoparticles are selected from at least one of Fe-Co alloy, Fe-Co-V alloy and Fe-Ni alloy, and the particle size of the soft magnetic alloy nanoparticles is from 50 nm to 200 nm.

4. The multi-component composite alloy additive according to claim 3, characterized in that, The soft magnetic alloy nanoparticles are Fe-Co-V alloys, and by mass percentage, the Fe-Co-V alloy contains 50% Fe, 48% Co, and 2% V.

5. The multi-component composite alloy additive according to claim 1 or 2, characterized in that, The transient liquid phase alloy is a Ni-Cu-P alloy, wherein, by mass percentage, P is 9% to 12%, Cu is 5% to 15%, and the balance is Ni and unavoidable impurities.

6. The multi-component composite alloy additive according to claim 1 or 2, characterized in that, The two-dimensional layered ceramic nanosheets are hexagonal boron nitride nanosheets, with a thickness of no more than 5 nm and a lateral diameter of 200 nm to 500 nm.

7. A method for preparing a multi-component composite alloy additive as described in any one of claims 1 to 6, characterized in that, include: The soft magnetic alloy nanoparticles are sensitized and activated, then dispersed in a chemical plating solution containing nickel salt, copper salt, and hypophosphite for plating. A transient liquid phase alloy shell is deposited on the outer surface of the soft magnetic alloy nanoparticles to obtain coated particles. The coated particles are then dry-mixed with the two-dimensional layered ceramic nanosheets in an inert atmosphere, so that the two-dimensional layered ceramic nanosheets adhere to the outer surface of the coated particles.

8. The preparation method according to claim 7, characterized in that, The electroless plating solution has a pH of 4.5 to 5.5 and a temperature of 80°C to 90°C; the dry mixing speed is 150 r / min to 250 r / min, and the net running time is 30 min to 60 min.

9. The application of the multi-component composite alloy additive as described in any one of claims 1 to 6 in the preparation of porous and breathable plastic mold steel, characterized in that, The application includes mixing the multi-component composite alloy additive with plastic mold steel matrix powder, molding, and performing stepped heating sintering. The amount of the multi-component composite alloy additive is 5% to 15% of the mass of the plastic mold steel matrix powder. The stepped heating sintering includes a first holding stage and a second holding stage. The temperature of the first holding stage is higher than the liquidus temperature of the transient liquid phase alloy and lower than the solidus temperature of the plastic mold steel matrix powder. The temperature of the second holding stage is higher than the temperature of the first holding stage and lower than the lower of the solidus temperature of the plastic mold steel matrix powder and the melting point of the metal element with the highest mass fraction in the transient liquid phase alloy.

10. The application according to claim 9, characterized in that, The plastic mold steel matrix powder is martensitic stainless steel powder or hot work mold steel powder with a particle size of 15 μm to 45 μm. The molding process is cold isostatic pressing, with a pressure of 150 MPa to 250 MPa and a holding time of 5 min. The stepped heating sintering includes a degreasing stage in which the temperature is increased to 430°C to 480°C at a rate of 1°C to 3°C per minute in a hydrogen-containing reducing atmosphere and held for 2.5 h to 3.5 h; a first holding stage in which the temperature is increased to 930°C to 960°C at a rate of 2°C to 4°C per minute in a vacuum not exceeding 10⁻² Pa and held for 1.5 h to 3 h; and a second holding stage in which the temperature is increased to 1150°C to 1170°C at a rate of 4°C to 6°C per minute in an inert gas partial pressure of 200 Pa to 1000 Pa and held for 2.5 h to 3.5 h. The mass fraction of the melting point reducing element in the obtained porous and breathable plastic mold steel is 0.3% to 0.8%.

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