Nanometer alumina dispersion strengthened iron-based composite powder for powder metallurgy gear, preparation method and powder metallurgy gear

Abstract

This invention belongs to the field of metal dispersion strengthening and powder metallurgy technology, specifically relating to nano-alumina dispersion-strengthened iron-based composite powder for powder metallurgy gears, its preparation method, and powder metallurgy gears. The composite powder has α-Fe as the continuous matrix phase, wherein dispersedly distributed iron-aluminum spinel Fe with α-Al₂O₃ as the core and magnesium as the stabilizer. (1‑x) Mg x Al2O4 is a core-shell type oxide dispersion with an interface shell. The interface shell continuously coats the core surface and forms a semi-coherent interface with lattice mismatch of no more than 6% with α-Fe. During preparation, iron nitrate, aluminum nitrate and magnesium salt are co-precipitated to obtain a precursor, which is then thermally decomposed in situ to generate a core-shell oxide. The oxide is then reduced with hydrogen under a controlled reduction window that reduces the iron oxide while retaining the interface shell. This powder transforms the alumina dispersion particles from weak chains that induce fatigue cracks into a well-bonded reinforcement that resists cyclic dissociation, significantly improving the high-temperature bending fatigue limit of powder metallurgy gears.

Description

Technical Field

[0001] This invention belongs to the field of metal dispersion strengthening and powder metallurgy technology, specifically relating to nano-alumina dispersion strengthened iron-based composite powder for powder metallurgy gears, preparation method, and powder metallurgy gears. Background Technology

[0002] Powder metallurgy gears are widely used in transmission components such as automotive gearboxes, power tools, and industrial speed reducers due to their advantages of near-net-shape forming, high material utilization, and good batch consistency. As transmission systems develop towards higher power density, the service oil temperature of gears continues to rise. In automotive automatic transmissions, the oil temperature often exceeds 150 °C under frequent shifting and heavy-load conditions. The grain boundary strength of iron-based sintered bodies decreases rapidly with increasing temperature, resulting in insufficient fatigue resistance of gears at high temperatures. Bending fatigue cracks at the tooth root are prone to initiation and propagation under cyclic loading, becoming a key bottleneck restricting the high-temperature, heavy-load application of powder metallurgy gears.

[0003] To improve the high-temperature performance of iron-based materials, the field typically employs two main approaches: alloying strengthening and oxide dispersion strengthening. Alloying strengthening relies on solid solution strengthening and precipitation strengthening caused by solid solution elements in the matrix. However, the diffusion of solid solution elements accelerates at high temperatures, easily leading to grain growth and coarsening of precipitates, causing the material to soften during long-term high-temperature service and making it difficult to maintain high-temperature fatigue resistance. Oxide dispersion strengthening, on the other hand, introduces thermally stable nano-oxide second-phase particles into the metal matrix, pinning dislocations and grain boundaries and hindering dislocation movement, thus maintaining strength and creep resistance at high temperatures. It is a mature strengthening method in the fields of high-performance copper alloys, high-temperature alloys, and high-chromium ferritic steels.

[0004] In the preparation of oxide dispersion-reinforced iron-based raw material powders, traditional methods mainly include mechanical alloying, internal oxidation, and reactive gas atomization. Mechanical alloying forcibly mixes nano-oxides with pre-alloyed powders through prolonged ball milling. However, the oxides are difficult to disperse uniformly during ball milling and tend to agglomerate, resulting in coarse and unevenly distributed dispersed phase particles. Furthermore, a significant amount of oxide residue remains on the powder surface, deteriorating formability. Internal oxidation and reactive gas atomization form an oxide layer on the powder surface or a dispersed phase within the matrix, but they also suffer from the common problems of coarse and unevenly distributed oxides, making it difficult for the mechanical properties of subsequent thermoformed products to meet engineering application requirements.

[0005] To overcome the agglomeration of nanoparticles caused by uneven mechanical mixing, a wet chemical coprecipitation route has been proposed in this field. For example, a method for preparing nanoscale alumina-dispersed iron powder (patent number CN100425376C, invention title "Method for Preparing Nanoscale Alumina-Dispersed Iron Powder") discloses a method of preparing a mixed salt solution of ferric nitrate and aluminum nitrate, adding ammonia to coprecipitate ferric hydroxide and aluminum hydroxide, utilizing the opposite charges of ferric hydroxide colloid and aluminum hydroxide colloid near their isoelectric points to achieve uniform mixing at the nanoscale, followed by washing, heating decomposition, and hydrogen reduction to obtain dispersed iron powder with alumina particle diameter of 10-50 nm, and indicating that the dispersed iron powder is suitable for iron-based powder products requiring good high-temperature performance. This method achieves uniform dispersion of nano-alumina in an iron matrix through wet chemical coprecipitation, improving formability and reducing the probability of micron-sized coarse oxides.

[0006] However, the inventors discovered in-depth research that the dispersed iron powder prepared by the aforementioned wet chemical coprecipitation route has a reinforcing phase of pure α-Al₂O₃, and the interface between α-Al₂O₃ and the α-Fe matrix is ​​an incoherent interface with high interfacial energy and weak bonding. Under high-temperature cyclic thermomechanical loading at 150°C, the interface between the incoherent oxide and the iron matrix becomes a preferred site for interfacial dissociation and fatigue microcrack initiation; simultaneously, the inert oxide nanoparticles are easily sheared and dissolved in the slip zone of cyclic loading, leading to cyclic softening. In other words, although the pure α-Al₂O₃ dispersed particles are beneficial for monotonic strength and hardness, their interface with the matrix becomes a weak link in fatigue performance under cyclic loading. The beneficial effects addressed by the aforementioned patent focus on strength, hardness, and wear resistance, but do not solve the deeper technical problem of insufficient bending fatigue limit caused by the incoherence of the oxide-matrix interface under high-temperature cyclic loading.

[0007] Furthermore, a key characteristic distinguishing powder metallurgy gears from dense forged gears is the unavoidable residual porosity within their sintered bodies. These residual pores on the surface and subsurface of the sintered body often become preferential initiation sources for bending fatigue cracks at the tooth root. The size, morphology, and aggregation state of these pores at grain boundaries significantly influence the initiation and early propagation of fatigue cracks. Under high-temperature, heavy-load cyclic conditions, the stress concentration at residual pores and the interfacial dissociation at oxide interfaces superimpose, further exacerbating the decline in the high-temperature bending fatigue limit of powder metallurgy gears. In the mechanical alloying route, even with hydrogen-air ball milling to refine the oxides, nano-oxides still tend to agglomerate at grain boundaries, concentrating fatigue damage in weak areas lacking oxides. In the internal oxidation route, oxygen is difficult to uniformly penetrate the matrix, and the reaction ratio with the solute to be oxidized is difficult to precisely control, resulting in an unevenly distributed dispersed phase. These common limitations of existing raw material powders consistently restrict the ability of powder metallurgy gears to balance high-temperature bending strength and high-temperature bending fatigue limit.

[0008] Therefore, how to eliminate the weak chains of fatigue initiation under high-temperature cyclic loading at the level of the interface between the oxide and the iron matrix while maintaining the uniform dispersion of nano-alumina, thereby significantly improving the high-temperature bending fatigue limit of powder metallurgy gears, is a technical problem that urgently needs to be solved in this field. Summary of the Invention

[0009] The purpose of this invention is to overcome the defects in the prior art where the interface between the pure phase alumina dispersed particles and the iron matrix is ​​non-coherent and becomes a weak chain for fatigue crack initiation under high-temperature cyclic loading. This invention provides a nano-alumina dispersed reinforced iron-based composite powder for powder metallurgy gears, a preparation method, and powder metallurgy gears. The invention transforms the alumina dispersed particles from a weak chain for fatigue crack initiation into a well-bonded reinforcement that resists cyclic dissociation. This significantly improves the bending strength retention rate and bending fatigue limit of powder metallurgy gears at 150 °C while maintaining the uniform dispersion and grain refinement of nano-alumina.

[0010] To achieve the above objectives, the technical solution adopted by the present invention is summarized as follows.

[0011] This invention provides a nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears. The composite powder comprises an α-Fe continuous matrix phase and core-shell oxide dispersion particles dispersed within the α-Fe continuous matrix phase. The core-shell oxide dispersion particles include an α-Al₂O₃ core and a magnesium-stabilized iron-aluminum spinel interface layer as the shell. The magnesium-stabilized iron-aluminum spinel interface layer is composed of Fe. (1-x) Mg x Al₂O₄ with x of 0.1~0.5 is continuously coated on the surface of the α-Al₂O₃ core by in-situ reaction generated during thermal decomposition. The magnesium-stabilized iron-aluminum spinel interface layer and the α-Fe continuous matrix phase form a semi-coherent interface with a lattice mismatch of no more than 6%. Based on the total mass of the composite powder, the mass fraction of α-Fe is not less than 98%, the mass fraction of α-Al₂O₃ is 0.45%~1.79%, and the magnesium content is 0.01%~0.4% based on the magnesium-iron molar ratio; the particle size of the core-shell oxide dispersed particles is 10~30 nm.

[0012] Preferably, the thickness of the magnesium-stabilized iron-aluminum spinel interface layer is 1~3 nm, and it satisfies the orientation relationship of the spinel {111} plane being parallel to the α-Fe {110} plane with the α-Fe continuous matrix phase; the Fe (1-x) Mg xThe x content in Al2O4 is further 0.25~0.35. The interface layer maintains its spinel structure without decomposition at a sintering temperature of 1200 °C via a magnetite-type inverse spinel mechanism. The volume fraction of the interface layer in the total α-Al2O3 in the core-shell oxide dispersed particles is 5%~20%. As a further preferred embodiment, the composite powder also contains chromium, which partially replaces the aluminum in the iron-aluminum spinel interface layer to form a (Fe,Mg)(Al,Cr)2O4 solid solution interface layer and is simultaneously dissolved in the α-Fe continuous matrix phase. The chromium content is 0.1%~1.0% based on the chromium-iron molar ratio.

[0013] This invention also provides a method for preparing the above-mentioned iron-based composite powder, comprising the following steps: Step 1, dissolving ferric nitrate, aluminum nitrate, and magnesium salt in deionized water to prepare a mixed salt solution, and co-precipitating with ammonia water to simultaneously precipitate ferric hydroxide, aluminum hydroxide, and magnesium hydroxide to obtain an iron-aluminum-magnesium mixed hydroxide precursor, with the pH at the coprecipitation endpoint controlled at 9.5~10.0. Step 2, after washing the precursor, thermally decomposing it at a temperature that allows iron oxide and aluminum oxide to react at the interface to form the magnesium-stabilized iron-aluminum spinel interface layer, to obtain an intermediate product in an iron oxide matrix, with α-Al2O3 as the core and the magnesium-stabilized iron-aluminum spinel interface layer as the shell. Step 3, subjecting the intermediate product to hydrogen reduction within a controlled reduction window that reduces iron oxide to α-Fe while retaining the magnesium-stabilized iron-aluminum spinel interface layer, to obtain the iron-based composite powder. Preferably, the controlled reduction window in step three is a reduction temperature of 480~600 °C, a reduction gas dew point of -40~-20 °C, and a reduction time of 30~60 min; the thermal decomposition temperature in step two is 800~900 °C, and the precursor is sent into the thermal decomposition furnace after ultrasonic spraying for thermal decomposition.

[0014] The present invention also provides a powder metallurgy gear, which is made by pressing and sintering the aforementioned iron-based composite powder. The powder metallurgy gear retains no less than 87% of its room temperature value in bending strength at 150 °C.

[0015] Compared with the prior art, the beneficial effects of the present invention are as follows. First, the present invention introduces trace amounts of magnesium simultaneously during the co-precipitation stage and terminates hydrogen reduction at a controlled reduction window that both reduces bulk iron oxide to α-Fe and retains the magnesium-stabilized iron-aluminum spinel interface layer at the oxide-metal interface, thereby forming a continuously coated semi-coherent spinel interface shell in situ on the surface of the α-Al2O3 core. The interface shell has a low mismatch with the α-Al2O3 core due to sharing a close-packed oxygen sublattice, and a low mismatch with the α-Fe matrix due to the orientation relationship between the spinel {111} plane and the α-Fe {110} plane. Thus, the original non-coherent α-Al2O3 and α-Fe interface with a mismatch greater than 15% is transformed into a semi-coherent interface with a mismatch of no more than 6% in each sub-interface in a two-step bridging manner. This gives the interface mechanical compliance, resistance to cyclic dissociation, and resistance to shear dissolution, transforming the alumina dispersed particles from weak chains of fatigue crack initiation into a well-bonded reinforcement resistant to cyclic dissociation. Secondly, the introduced magnesium has multiple functions: it stabilizes the iron-aluminum spinel interface layer through a magnetite-type inverse spinel mechanism, allowing it to pass through sintering at 1200 °C without decomposition, reducing interfacial energy, and assisting in grain refinement and sintering densification. This demonstrates a synergistic effect of adding components to solve multiple deep technical problems simultaneously. Thirdly, this invention can significantly refine the grain size of the composite powder after sintering compared to ordinary iron powder, and ensures that the bending strength retention rate of powder metallurgy gears at 150 °C is not less than 87% of its room temperature value. Furthermore, it significantly improves the high-temperature bending fatigue limit compared to pure phase alumina dispersed iron powder, providing a high-temperature resistant raw material powder and its co-precipitation-controlled reduction preparation scheme for high-temperature heavy-duty powder metallurgy gears.

[0016] The core-shell alumina dispersion-reinforced iron-based composite powder provided by this invention retains the advantages of the wet chemical co-precipitation route in terms of uniform dispersion and good formability of nano-alumina, while fundamentally improving the interfacial bonding between the oxide and the iron matrix through interface engineering. Therefore, it is particularly suitable for powder metallurgy gears operating under high-temperature, heavy-load cyclic conditions, such as automotive automatic transmission gears, power tool gears, and industrial reducer gears. The composite powder can be pressed and sintered using conventional powder metallurgy processes, without the need for complex large-scale equipment or harsh operating conditions. The process control is simple, the cost is low, and it is environmentally friendly, facilitating large-scale production. Furthermore, by adjusting parameters such as the aluminum-iron molar ratio, magnesium substitution degree, interface shell integral number, and whether chromium is introduced, the dispersion reinforcement degree, interface matching degree, and high-temperature oxidation resistance of the matrix can be controlled within a wide range to meet the differentiated requirements of strength, toughness, and high-temperature bending fatigue limit for various high-temperature, heavy-load gears. Attached Figure Description

[0017] Figure 1 This is a process flow diagram of the preparation process of nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears in Embodiment 1 of the present invention.

[0018] Figure 2 This is a graph showing the relationship between the relative reduction rate of the magnesium-stabilized iron-aluminum spinel interface layer and temperature within the controlled reduction window of this invention.

[0019] Figure 3 This is a comparison chart of the bending fatigue stress-life curves of the powder metallurgy gears of Example 1 and Comparative Example 1 at 150 °C. Detailed Implementation

[0020] To make the objectives, technical solutions, and beneficial effects of this invention clearer, the invention will be further described in detail below with reference to the accompanying drawings and specific embodiments. It should be understood that the specific embodiments described herein are only for explaining the invention and are not intended to limit the invention. The raw materials used in the following embodiments, including ferric nitrate, aluminum nitrate, magnesium nitrate, and chromium nitrate, are all analytically pure, calculated as ferric nitrate nonahydrate Fe(NO3)3·9H2O, aluminum nitrate nonahydrate Al(NO3)3·9H2O, magnesium nitrate hexahydrate Mg(NO3)2·6H2O, and chromium nitrate nonahydrate Cr(NO3)3·9H2O; the ammonia water mass fraction is 25%; and the hydrogen purity is not less than 99.99%.

[0021] In the various embodiments of this invention, the preparation of the mixed salt solution, precise control of the pH at the co-precipitation endpoint, the atomization frequency and carrier gas flow rate of the ultrasonic spray, the temperature of the pyrolysis furnace and the residence time of the precursor in the furnace, and the temperature, reducing gas dew point, and reduction time during the hydrogen reduction process are all key process parameters affecting the core-shell structure and performance of the final composite powder. Specifically, the pH at the co-precipitation endpoint determines whether the iron, aluminum, and magnesium phases can precipitate completely and simultaneously; the pyrolysis temperature determines whether iron oxide and aluminum oxide can react at the interface to form a magnesium-stabilized iron-aluminum spinel interface layer; and the temperature of the controlled reduction window and the reducing gas dew point jointly determine whether the magnesium-stabilized iron-aluminum spinel interface layer can be retained while the bulk iron oxide is reduced to α-Fe. The following embodiments prepare composite powders under different values ​​of the above key process parameters, and their core-shell structural characteristics and high-temperature mechanical properties are tested.

[0022] The overall preparation process of this invention is as follows: Figure 1 As shown, the process sequentially includes five stages: preparation of a mixed salt solution, co-precipitation with ammonia and control of the final pH, washing and ultrasonic spraying, in-situ thermal decomposition at 800-900 °C to generate core-shell oxides, and hydrogen reduction under a controlled reduction window. The final product is an iron-based composite powder with α-Fe as the continuous matrix phase and dispersed core-shell oxide particles. The dispersed core-shell oxide particles, from the inside out, consist of an α-Al₂O₃ core and a continuously coated magnesium-stabilized iron-aluminum spinel interface layer. This interface layer and the outer α-Fe continuous matrix phase form a semi-coherent interface.

[0023] Example 1

[0024] This embodiment prepares a core-shell alumina dispersion-reinforced iron-based composite powder with an aluminum-iron molar ratio of 1.25% and a magnesium-iron molar ratio of 0.1%. 202 g of ferric nitrate nonahydrate, 2.34 g of aluminum nitrate nonahydrate, and 0.13 g of magnesium nitrate hexahydrate were weighed and dissolved in deionized water to prepare a 250 mL mixed salt solution, which was stirred until completely dissolved. While continuously stirring, 25% ammonia solution was added to the mixed salt solution at a rate of 10 mL / s to simultaneously precipitate ferric hydroxide, aluminum hydroxide, and magnesium hydroxide. After the addition was complete, ammonia solution was continued to be added dropwise to precisely adjust the pH of the coprecipitation endpoint to 9.75, and stirring was continued at this pH for 2 min to ensure simultaneous and complete precipitation of the three phases. Since magnesium hydroxide has a relatively small solubility product, controlling the pH of the coprecipitation endpoint between 9.5 and 10.0 is necessary to ensure the simultaneous and complete coprecipitation of magnesium with iron and aluminum, thereby providing magnesium as a stabilizer in the subsequent interfacial shell.

[0025] The resulting coprecipitate was added to 1 L of deionized water, centrifuged, and washed four times to remove residual nitrate and ammonium ions, yielding a slurry of iron-aluminum-magnesium mixed hydroxide precursors. The precursor slurry was atomized into fine droplets using ultrasonic spraying and fed into a thermal decomposition furnace at 850 °C for thermal decomposition. During this thermal decomposition, the hydroxides dehydrated and decomposed to generate iron oxide, α-Al₂O₃, and magnesium oxide. An interfacial reaction occurred at the interface between iron oxide and α-Al₂O₃, resulting in the in-situ formation of magnesium-stabilized iron-aluminum spinel Fe, continuously coated on the surface of the α-Al₂O₃ core. (1-x) Mg x An Al2O4 interface layer was obtained, yielding an intermediate product in which core-shell oxides were dispersed in an iron oxide matrix.

[0026] The intermediate product was subjected to hydrogen reduction under a controlled reduction window at a reduction temperature of 540 °C, a reducing gas dew point of -30 °C, and a reduction time of 45 min. Within this controlled reduction window, bulk iron oxide was rapidly reduced to α-Fe, while the magnesium-stabilized iron-aluminum spinel interface layer was preserved due to its high activation energy required for reduction and the freezing of reduction kinetics at this temperature. This resulted in an iron-based composite powder with α-Fe as the continuous matrix phase and dispersed α-Al₂O₃ cores and a magnesium-stabilized iron-aluminum spinel interface layer as the shell. Transmission electron microscopy and energy dispersive spectroscopy analysis showed that the particle size of the core-shell oxide dispersions was approximately 20 nm, the thickness of the interface shell was approximately 2 nm, the x-coefficient of the interface shell composition was approximately 0.3, and the volume fraction of the interface shell in the total α-Al₂O₃ was approximately 12.5%. Selected area electron diffraction showed that the interface shell and the α-Fe matrix had an orientation relationship where the spinel {111} plane was parallel to the α-Fe {110} plane, forming a semi-coherent interface. Based on the total mass of the composite powder, the mass fraction of α-Al₂O₃ was approximately 1.13%, and the mass fraction of α-Fe was not less than 98%.

[0027] The iron-based composite powder obtained in this embodiment was further characterized. X-ray diffraction analysis showed that the phase of the composite powder consisted of α-Fe and a small amount of α-Al₂O₃ and Fe. (1-x) Mg x The composition is Al2O4, and no diffraction peaks of free magnesium oxide were detected, indicating that magnesium mainly enters the iron-aluminum spinel interface layer. Specific surface area measurements show that the composite powder has a specific surface area of ​​3.2 m². 2 / g, oxygen content is 0.61%, and loose density is 2.85 g / cm³. 3 The tap density is 3.42 g / cm³. 3 The flowability, measured by a Hall effect flowmeter, is 28 s / 50 g, and the compressibility at a pressing pressure of 600 MPa corresponds to a compact density of 7.05 g / cm³. 3 Because the magnesium-stabilized iron-aluminum spinel interface layer encapsulates the α-Al₂O₃ core during preparation, and there are fewer residual oxides on the α-Fe matrix surface, the resulting composite powder exhibits good formability and compressibility. High-angle annular dark-field images and energy dispersive spectroscopy further confirm that aluminum and magnesium elements are concentrated in the core-shell region of the dispersed particles, and the magnesium signal is stronger at the interface shell than in the core region, consistent with the core-shell structure of the magnesium-stabilized iron-aluminum spinel interface layer continuously encapsulating the α-Al₂O₃ core surface.

[0028] The obtained iron-based composite powder was pressed into gear blanks under a pressure of 800 MPa and sintered at 1200 °C for 2 h to obtain powder metallurgy gears. Testing showed that the sintered gears produced in this example had a grain size of approximately 2.6 μm, a room temperature flexural strength of 980 MPa, a flexural strength at 150 °C of 892 MPa, and a flexural strength retention rate of 91.0% at 150 °C. 7 The flexural fatigue limit measured in the second cycle was 342 MPa, and the sintering density was 7.52 g / cm³. 3 .

[0029] Example 2

[0030] This embodiment prepares a core-shell alumina dispersion-reinforced iron-based composite powder with an aluminum-iron molar ratio of 0.5% and a magnesium-iron molar ratio of 0.04%, corresponding to a lower limit scheme of approximately 0.45% α-Al₂O₃ mass fraction. 202 g of ferric nitrate nonahydrate, 0.94 g of aluminum nitrate nonahydrate, and 0.05 g of magnesium nitrate hexahydrate were weighed and dissolved in deionized water to prepare a 250 mL mixed salt solution. Ammonia solution with a mass fraction of 25% was added at a rate of 10 mL / s under continuous stirring to induce simultaneous precipitation of the three phases. The pH of the coprecipitation endpoint was precisely adjusted to 9.5, and stirring was continued for 2 min. The coprecipitate was washed four times to obtain a precursor slurry, which was then ultrasonically sprayed and sent to an 800 °C thermal decomposition furnace for pyrolysis, generating a core-shell oxide intermediate in situ. The intermediate was then reduced with hydrogen under a controlled reduction window of 480 °C, a reducing gas dew point of -40 °C, and a reduction time of 30 min to obtain the iron-based composite powder.

[0031] X-ray diffraction analysis showed that the phase composition of the composite powder in this embodiment consisted of α-Fe and a small amount of α-Al₂O₃ and Fe. (1-x) Mg x Composed of Al2O4, the composite powder has a specific surface area of ​​2.9 m². 2 / g, oxygen content is 0.42%, and loose density is 2.92 g / cm³. 3 Because the aluminum-iron molar ratio in this embodiment is at the lower limit, the size of the α-Al₂O₃ core and its interface shell is correspondingly smaller.

[0032] Testing revealed that the core-shell oxide dispersions prepared in this embodiment had a particle size of approximately 11 nm, an interface shell thickness of approximately 1.1 nm, and an interface shell composition corresponding to x of approximately 0.25. The interface shell and the α-Fe matrix formed a semi-coherent interface. The resulting composite powder was pressed and sintered under the same conditions as in Example 1 to produce a powder metallurgy gear. After sintering, the gear had a grain size of approximately 3.0 μm, a room temperature flexural strength of 920 MPa, a flexural strength at 150 °C of 814 MPa, a flexural strength retention rate of 88.5% at 150 °C, a flexural fatigue limit of 315 MPa at 150 °C, and a sintering density of 7.48 g / cm³. 3 .

[0033] Example 3

[0034] This embodiment prepares a chromium-containing core-shell alumina dispersion-reinforced iron-based composite powder, corresponding to the upper limit of the aluminum-iron molar ratio of 2.0%. Chromium is introduced to further adjust interfacial mismatch, improve the high-temperature oxidation resistance of the matrix, and promote densification. 202 g of ferric nitrate nonahydrate, 3.75 g of aluminum nitrate nonahydrate, 0.50 g of magnesium nitrate hexahydrate, and 1.00 g of chromium nitrate nonahydrate were weighed and dissolved in deionized water to prepare a 250 mL mixed salt solution, making the chromium-iron molar ratio 0.5%. Ammonia solution with a mass fraction of 25% was added at a rate of 10 mL / s under continuous stirring, causing ferric hydroxide, aluminum hydroxide, magnesium hydroxide, and chromium hydroxide to precipitate simultaneously. The pH of the co-precipitation endpoint was precisely adjusted to 10.0, and stirring was continued for 2 min. The precursor slurry was obtained by co-precipitation and washing four times. After ultrasonic spraying, it was sent to a 900 °C pyrolysis furnace for in-situ decomposition to generate a core-shell oxide intermediate with an α-Al₂O₃ core and a (Fe,Mg)(Al,Cr)₂O₄ solid solution interface layer as the shell. The intermediate was then reduced with hydrogen at a controlled reduction window of 600 °C, a reducing gas dew point of -20 °C, and a reduction time of 60 min to obtain a chromium-iron-based composite powder. The chromium portion replaced the aluminum in the interface layer to form the (Fe,Mg)(Al,Cr)₂O₄ solid solution interface layer, and simultaneously dissolved in the α-Fe continuous matrix phase.

[0035] Energy dispersive spectroscopy (EDS) scans showed that chromium was simultaneously distributed in both the dispersed particle interface shell and the α-Fe matrix in this embodiment: at the interface shell, chromium coexisted with aluminum, corresponding to chromium partially replacing aluminum in the interface layer to form a (Fe,Mg)(Al,Cr)₂O₄ solid solution interface layer; a dispersed chromium signal was detected in the matrix, corresponding to chromium solid dissolved in the continuous α-Fe matrix phase. X-ray diffraction did not detect independent chromium oxide or chromium intermetallic compound diffraction peaks. The specific surface area of ​​the composite powder was 3.5 m². 2 / g, with an oxygen content of 0.78%.

[0036] Testing revealed that the core-shell oxide dispersion particles prepared in this embodiment had a particle size of approximately 29 nm, an interface shell thickness of approximately 2.9 nm, and an interface shell composition corresponding to x of approximately 0.35. The resulting composite powder was pressed and sintered under the same conditions as in Example 1 to produce a powder metallurgy gear. The sintered gear had a grain size of approximately 2.4 μm, a room temperature flexural strength of 1015 MPa, a flexural strength at 150 °C of 928 MPa, a flexural strength retention rate at 150 °C of 91.4%, a flexural fatigue limit at 150 °C of 356 MPa, and a sintering density of 7.55 g / cm³. 3 It is evident that the introduction of chromium, while maintaining the retention rate of high-temperature bending strength, further improved the bending fatigue limit and increased the sintering density, verifying the beneficial role of chromium as a synergistic component that simultaneously regulates interfacial mismatch, enhances the high-temperature oxidation resistance of the matrix, and promotes densification.

[0037] Example 4

[0038] This embodiment investigates the effect of magnesium salt type on core-shell structure and performance. Magnesium acetate was used instead of magnesium nitrate as the magnesium source, and the remaining preparations were carried out using the same aluminum-iron molar ratio of 1.25% and magnesium-iron molar ratio of 0.1% as in Example 1. 202 g of ferric nitrate nonahydrate, 2.34 g of aluminum nitrate nonahydrate, and 0.11 g of magnesium acetate tetrahydrate were weighed and dissolved in deionized water to prepare a 250 mL mixed salt solution. Ammonia solution with a mass fraction of 25% was added at a rate of 10 mL / s under continuous stirring to simultaneously precipitate ferric hydroxide, aluminum hydroxide, and magnesium hydroxide. The pH of the coprecipitation endpoint was precisely adjusted to 9.75, and stirring was continued for 2 min. The coprecipitate was washed four times to obtain a precursor slurry, which was then ultrasonically sprayed and sent to an 850 °C thermal decomposition furnace for pyrolysis, generating a core-shell oxide intermediate in situ. The intermediate was reduced with hydrogen under a controlled reduction window of 540 °C, a reducing gas dew point of -30 °C, and a reduction time of 45 min to obtain an iron-based composite powder. Testing revealed that the particle size of the core-shell oxide dispersions prepared in this embodiment was approximately 20 nm, the interface shell thickness was approximately 2 nm, and the x-value corresponding to the interface shell composition was approximately 0.3, which is basically consistent with Example 1. This indicates that, under the premise of ensuring simultaneous co-precipitation of magnesium with iron and aluminum, the type of magnesium source has no significant impact on the core-shell structure. The obtained composite powder was pressed and sintered under the same conditions as in Example 1 to prepare powder metallurgy gears. The gears exhibited a room temperature bending strength of 975 MPa, a bending strength at 150 °C of 886 MPa, a bending strength retention rate of 90.9% at 150 °C, a bending fatigue limit of 339 MPa at 150 °C, and a sintering density of 7.51 g / cm³. 3 .

[0039] Example 5

[0040] This embodiment investigates the effect of the interface shell integral on performance. By appropriately extending the thermal decomposition time, the interfacial reaction is made more complete, thereby increasing the volume fraction of the interface shell in the total α-Al₂O₃ to approximately 18%, close to the upper limit, while maintaining the aluminum-iron molar ratio of 1.25%. A mixed salt solution was prepared according to the same proportions as in Example 1, and after co-precipitation, washing, and ultrasonic spraying, thermal decomposition was performed at 870°C. The residence time of the precursor in the furnace was appropriately extended to allow for a more complete reaction between iron oxide and aluminum oxide at the interface, generating an in-situ core-shell oxide intermediate with an interface shell integral of approximately 18%. This intermediate was then reduced with hydrogen at a controlled reduction window of 540°C, a dew point of -30°C, and 45 min to obtain an iron-based composite powder. The particle size of the dispersed particles of the core-shell oxide in this embodiment was approximately 21 nm, and the thickness of the interface shell was approximately 2.8 nm. The obtained composite powder was pressed and sintered under the same conditions as in Example 1 to prepare a powder metallurgy gear. Its room temperature flexural strength was 968 MPa, its flexural strength at 150 °C was 884 MPa, its flexural strength retention rate at 150 °C was 91.3%, its flexural fatigue limit at 150 °C was 345 MPa, and its sintering density was 7.50 g / cm³. 3 It can be seen that when the integral number of the interface shell is in the range of 5% to 20%, the high-temperature bending fatigue limit and strength retention rate are maintained at a high level. If the integral number of the interface shell is too high, it will consume α-Al2O3 cores and slightly weaken the dispersion strengthening. If it is too low, the interface will not be fully covered and incoherent weak regions will remain. Therefore, 5% to 20% is appropriate.

[0041] Comparative Example 1

[0042] This comparative example reproduces pure-phase alumina-dispersed iron powder using the existing wet chemical coprecipitation method, serving as a control to remove all the distinguishing technical features of this invention. 202 g of ferric nitrate nonahydrate and 2.34 g of aluminum nitrate nonahydrate were weighed, without the addition of magnesium nitrate, and dissolved in deionized water to prepare a 250 mL mixed salt solution. Ammonia solution with a mass fraction of 25% was added at a rate of 10 mL / s to coprecipitate the ferric hydroxide and aluminum hydroxide. The pH was adjusted to 9.0, and after washing four times, the solution was decomposed by heating at 200 °C, followed by reduction with hydrogen at 700 °C for 45 min to obtain pure-phase α-Al₂O₃ dispersed iron powder. Because it does not contain a magnesium stabilizer and the reduction temperature is higher than the controlled reduction window, the resulting dispersed particles are pure-phase α-Al₂O₃, with a non-coherent interface between the particle and the α-Fe matrix, lacking a magnesium-stabilized iron-aluminum spinel interface shell. The obtained powder was pressed and sintered under the same conditions as in Example 1 to produce a gear. The sintered gear had a grain size of approximately 2.7 μm, a room temperature flexural strength of 935 MPa, a flexural strength at 150 °C of 729 MPa, a 150 °C flexural strength retention rate of only 78.0%, a 150 °C flexural fatigue limit of only 258 MPa, and a sintered density of 7.41 g / cm³. 3 .

[0043] Comparative Example 2

[0044] This comparative example uses a mechanical mixing method to ball-mill and mix nano-α-Al₂O₃ powder and reduced iron powder, serving as a control to replace the in-situ core-shell particles generated in this invention with mechanical alloying characteristics closest to those of the prior art. Nano-α-Al₂O₃ powder with an average particle size of 30 nm and reduced iron powder were weighed at 1.13% by mass and ball-milled in a planetary ball mill for 48 h. Because nano-α-Al₂O₃ is difficult to disperse uniformly and easily agglomerates during ball milling, the resulting powder contains localized oxide agglomerates larger than 100 nm, and the interface between the oxides and the iron matrix is ​​incoherent. The resulting powder was pressed and sintered under the same conditions as in Example 1 to produce gears. The sintered grain size was approximately 4.5 μm, the room temperature bending strength was 840 MPa, the bending strength at 150 °C was 638 MPa, the 150 °C bending strength retention rate was 76.0%, the 150 °C bending fatigue limit was only 236 MPa, and the sintering density was 7.30 g / cm³. 3 The performance of the former is significantly inferior to that of the embodiments of the present invention, which shows the necessity of co-precipitation in-situ generation and interface engineering for achieving uniform dispersion and good interfacial bonding of nano-alumina.

[0045] Comparative Example 3

[0046] In this comparative example, aluminum nitrate was added during the co-precipitation stage, but magnesium nitrate was not added. The same controlled reduction window as in Example 1 was used as a control to remove the molecularly designed magnesium component and retain only the holmite interfacial shell. 202 g of ferric nitrate nonhydrate and 2.34 g of aluminum nitrate nonhydrate were weighed, without the addition of magnesium nitrate. The mixture was co-precipitated, washed, ultrasonically sprayed, and thermally decomposed at 850 °C in the same manner as in Example 1, generating a magnesium-free pure holmite FeAl2O4 interfacial shell in situ. This was then reduced with hydrogen at a controlled reduction window of 540 °C. Although the resulting powder formed a pure holmite interfacial shell after reduction, during the subsequent sintering process at 1200 °C, the magnesium-free pure holmite interfacial shell decomposed, losing its semi-coherent interface, and the interface reverted to an incoherent state. The obtained powder was pressed and sintered under the same conditions as in Example 1 to produce gears. The gears exhibited a room temperature flexural strength of 948 MPa, a flexural strength at 150 °C of 758 MPa, a flexural strength retention rate of 80.0% at 150 °C, a flexural fatigue limit of 271 MPa at 150 °C, and a sintered density of 7.43 g / cm³. 3 .

[0047] Comparative Example 4

[0048] This comparative example added magnesium nitrate during the co-precipitation stage but used a reduction temperature higher than the controlled reduction window, serving as a control for adding magnesium but removing the interfacial shell. Ferric nitrate, aluminum nitrate, and magnesium nitrate were added in the same proportions as in Example 1 for co-precipitation, washing, ultrasonic spraying, and thermal decomposition at 850 °C, followed by hydrogen reduction at 700 °C for 45 min. Because the reduction temperature was higher than the controlled reduction window, the magnesium-stabilized iron-aluminum spinel interfacial shell was further reduced and decomposed during the reduction process, and magnesium dissolved into the α-Fe matrix. The resulting powder contained pure α-Al₂O₃ oxide particles at a non-coherent interface with the matrix. The resulting powder was pressed and sintered under the same conditions as in Example 1 to produce gears. The gears exhibited a room temperature flexural strength of 942 MPa, a flexural strength at 150 °C of 754 MPa, a flexural strength retention rate of 80.0% at 150 °C, a flexural fatigue limit of 268 MPa at 150 °C, and a sintered density of 7.44 g / cm³. 3 .

[0049] Comparative Example 5

[0050] The only difference between this comparative example and Example 1 is that the reduction temperature exceeds the controlled reduction window, serving as a parameter boundary check to verify the critical significance of the controlled reduction window. The precursor was prepared and thermally decomposed according to the same proportions and process as in Example 1. Subsequently, the reduction temperature was increased from 540 °C to 700 °C, while other conditions remained unchanged, for hydrogen reduction. At 700 °C, the reduction kinetics of the magnesium-stabilized iron-aluminum spinel interface shell no longer froze and decomposed, reducing the interface shell to α-Fe and α-Al₂O₃. The oxide dispersions in the resulting powder reverted to pure α-Al₂O₃, forming a non-coherent interface with the matrix. The resulting powder was pressed and sintered under the same conditions as in Example 1 to produce a gear. Its room temperature flexural strength was 930 MPa, its flexural strength at 150 °C was 726 MPa, its flexural strength retention at 150 °C was 78.1%, its flexural fatigue limit at 150 °C was 260 MPa, and its sintering density was 7.42 g / cm³. 3 This comparative example demonstrates that the controlled reduction window is a key process condition for preserving the semi-coherent interface shell and achieving the beneficial effects of this invention; exceeding this window will result in the decomposition of the interface shell and the loss of the beneficial effects.

[0051] The main instruments, equipment, and operating conditions used for performance testing in the above embodiments and comparative examples are as follows. Microstructure observation was performed using a field emission transmission electron microscope (FET) with an accelerating voltage of 200 kV. During sample preparation, the composite powder was dispersed in anhydrous ethanol via ultrasonic dispersion and then dropped onto a microgrid copper mesh. The interface shell thickness and core-shell size were measured in high-resolution mode. The chemical composition of the interface shell was collected at specific points using an attached energy dispersive spectroscopy (EDS) instrument in nanobeam mode, and the magnesium substitution degree (x) was calculated accordingly. The core-shell interface orientation relationship was determined using selected area electron diffraction combined with fast Fourier transform. Phase analysis was performed using an X-ray diffractometer with copper target Kα radiation, a scanning range of 20°–90°, and a scanning speed of 4° / min. Specific surface area was measured using a nitrogen adsorption specific surface area analyzer via a multi-point method. Oxygen content was measured using an oxygen, nitrogen, and hydrogen analyzer. Loose packing density and tapped density were measured using appropriate measuring instruments. Flowability was measured using a Hall effect flowmeter.

[0052] The testing methods for each performance item in the above embodiments and comparative examples are as follows. The particle size, interface shell thickness, and core-shell structure of the dispersed particles in the core-shell oxide were observed using high-resolution transmission electron microscopy. The chemical composition of the interface shell was quantitatively analyzed using energy dispersive spectroscopy (EDS) to calculate the magnesium substitution degree x. The orientation relationship between the interface shell and the matrix was determined using selected area electron diffraction (SID). The grain size after sintering was measured using a metallographic microscope using the truncation method. The room temperature bending strength and the 150 °C bending strength were determined using the three-point bending method according to the corresponding test standards. The 150 °C bending strength was measured after the sample temperature stabilized at 150 °C. The retention rate of the 150 °C bending strength was the ratio of the 150 °C bending strength to the room temperature bending strength. The 150 °C bending fatigue limit was determined using the rise-fall method at 150 °C with a stress ratio of 0.1, to withstand 10... 7 The maximum stress gauge that prevents fracture after one cycle. Sintered density was determined using the Archimedes displacement method. All performance data are averaged from five parallel samples.

[0053] The core-shell structure characteristics and mechanical property test results of Examples 1-5 and Comparative Examples 1-5 are summarized in Table 1 for intuitive comparison, as shown in Table 1.

[0054] Table 1. Core-shell structural features and mechanical property test results of the examples and comparative examples.

[0055] sample Particle size / nm Interface shell thickness / nm Interface coherence Grain size / μm Room temperature bending strength / MPa 150 °C flexural strength / MPa Retention rate at 150 °C / % 150 °C fatigue limit / MPa Sintering density Example 1 20 2.0 semi-cognate 2.6 980 892 91.0 342 <![CDATA[7.52g / cm 3 ]]> Example 2 11 1.1 semi-cognate 3.0 920 814 88.5 315 <![CDATA[7.48g / cm 3 ]]> Example 3 29 2.9 semi-cognate 2.4 1015 928 91.4 356 <![CDATA[7.55g / cm 3 ]]> Example 4 20 2.0 semi-cognate 2.6 975 886 90.9 339 <![CDATA[7.51g / cm 3 ]]> Example 5 21 2.8 semi-cognate 2.5 968 884 91.3 345 <![CDATA[7.50g / cm 3 ]]> Comparative Example 1 22 0 Non-coherent 2.7 935 729 78.0 258 <![CDATA[7.41g / cm 3 ]]> Comparative Example 2 >100 reunions 0 Non-coherent 4.5 840 638 76.0 236 <![CDATA[7.30g / cm 3 ]]> Comparative Example 3 20 0 (after sintering) Non-coherent 2.8 948 758 80.0 271 <![CDATA[7.43g / cm 3 ]]> Comparative Example 4 21 0 Non-coherent 2.6 942 754 80.0 268 <![CDATA[7.44g / cm 3 ]]> Comparative Example 5 23 0 Non-coherent 2.9 930 726 78.1 260 <![CDATA[7.42g / cm 3 ]]>

[0056] Note: "0" for the interface shell thickness indicates that the dispersed particles are pure α-Al2O3 without magnesium-stabilized iron-aluminum spinel interface shell; "0 (after sintering)" in Comparative Example 3 indicates that a pure holmite interface shell exists after reduction but decomposes during sintering at 1200 °C.

[0057] As shown in Table 1, the embodiments of the present invention are significantly superior to the comparative examples in terms of high-temperature flexural strength retention rate and high-temperature flexural fatigue limit. Taking the optimal embodiment 1 as an example, its flexural strength retention rate at 150 °C reaches 91.0%, which is 13 percentage points higher than the 78.0% of Comparative Example 1, which replicates the existing wet chemical coprecipitation route; its flexural fatigue limit at 150 °C reaches 342 MPa, which is 32.6% higher than the 258 MPa of Comparative Example 1. This improvement does not come from the dispersion of nano-alumina and grain refinement itself, because Comparative Example 1 also achieved uniform dispersion of nano-alumina and similar grain size, but rather from the suppression of interface dissociation and shear dissolution under high-temperature cyclic loading by the magnesium-stabilized semi-coherent spinel interface shell constructed in situ on the surface of the α-Al2O3 core in the present invention.

[0058] Furthermore, the beneficial effects of this invention stem from the nonlinear synergy between the magnesium stabilizer and the semi-coherent interfacial shell, rather than a simple additive result of their contributions. Comparative Example 3, retaining the interfacial shell but removing the magnesium stabilizer, achieved a flexural fatigue limit of 271 MPa at 150 °C, an increase of only 13 MPa compared to the baseline of Comparative Example 1. Comparative Example 4, with magnesium added but the interfacial shell removed, achieved a flexural fatigue limit of 268 MPa at 150 °C, an increase of only 10 MPa compared to the baseline. If the two were simply additive, the expected fatigue limit would be 258 + 13 + 10, or 281 MPa; however, Example 1 achieved a measured value of 342 MPa, with a synergistic increase of 342 - 281, or 61 MPa, accounting for 73% of the total increase of 84 MPa in Example 1 relative to the baseline. The quantitative results indicate that the fatigue-inhibiting effect of the interface shell can only be fully realized when the magnesium stabilizer prevents the spinel interface shell from decomposing during sintering at 1200 °C, thus maintaining a semi-coherent interface in the final sintered body. In the absence of the magnesium stabilizer, the interface shell decomposes during sintering (Comparative Example 3), and in the absence of the interface shell, magnesium is only dissolved in the matrix while the interface remains incoherent (Comparative Example 4). Neither of these methods can individually achieve a significant improvement in fatigue, thus confirming that the combination of "magnesium stabilizer plus semi-coherent interface shell" produces a nonlinear synergistic effect far exceeding the sum of their individual contributions.

[0059] From the perspective of grain refinement, the grain size after sintering in all embodiments of the present invention is between 2.4 and 3.0 μm, which is significantly refined compared to the grain size of ordinary iron powder sintered bodies without the introduction of dispersed particles, with a refinement range of approximately 65%. This is attributed to the pinning effect of core-shell oxide dispersed particles on grain boundaries and the auxiliary refinement effect of magnesium. It should be noted that although Comparative Example 1 also introduced nano-alumina dispersed particles and obtained a grain size similar to that of the present invention, its 150 °C bending fatigue limit is significantly lower than that of the embodiments of the present invention. This indicates that grain refinement is not the main source of the improvement in high-temperature bending fatigue limit of the present invention, further confirming that the suppression of interface dissociation and shear dissolution by the semi-coherent interface shell is the key to fatigue improvement. From the perspective of strength retention rate, the 150 °C bending strength retention rate of Examples 1, 3, and 5 of the present invention all exceed 91%, while the retention rate of Comparative Example 3 and Comparative Example 4, which only use a single method, is 80%. The difference between the two again confirms the nonlinear synergy between magnesium stabilizer and semi-coherent interface shell.

[0060] The only difference between Comparative Example 5 and Example 1 is that the reduction temperature was increased from 540 °C within the controlled reduction window to 700 °C outside the window. Its 150 °C flexural fatigue limit plummeted from 342 MPa to 260 MPa, almost returning to the baseline level of Comparative Example 1, directly demonstrating the critical significance of the controlled reduction window for preserving the semi-coherent interface shell and achieving the beneficial effects of this invention. Comparative Example 2 used mechanical mixing; due to the agglomeration of nano-alumina and the incoherence of the interface, its high-temperature performance was the worst, further highlighting the necessity of combining in-situ co-precipitation generation with interface engineering in this invention. Chromium-containing Example 3 further increased the 150 °C flexural fatigue limit to 356 MPa and the sintering density to 7.55 g / cm³, based on Example 1. 3 This study validates the synergistic value of chromium as a single component that simultaneously regulates interfacial mismatch, enhances the high-temperature oxidation resistance of the matrix, and promotes densification.

[0061] The core mechanism of this invention lies in reconstructing the interface between alumina dispersed particles and the iron matrix through an in-situ constructed magnesium-stabilized semi-coherent spinel interface shell, thereby eliminating fatigue initiation weak chains under high-temperature cyclic loading. In existing pure-phase alumina dispersed iron powder, α-Al2O3 and α-Fe do not share crystallographic sublattices, and the interface lattice mismatch is greater than 15%, belonging to an incoherent high-interface-energy weak bonding interface. Under cyclic thermomechanical loading at 150 °C, this interface is prone to interfacial dissociation and becomes a preferential initiation site for fatigue microcracks. At the same time, the inert oxide is sheared and dissolved in the slip band, leading to cyclic softening.

[0062] This invention inserts a magnesium-stabilized iron-aluminum spinel interface shell between the α-Al₂O₃ core and the α-Fe matrix, transforming the originally single, highly mismatched interface into a two-step bridging semi-coherent interface cascade. On one hand, the spinel interface shell and the α-Al₂O₃ core share a close-packed oxygen sublattice, resulting in low mismatch: the oxygen-oxygen nearest neighbor spacing of the corundum basal plane is approximately 2.748 Å, while the oxygen-oxygen nearest neighbor spacing of the spinel oxygen sublattice is approximately 2.87~2.88 Å, with a mismatch of approximately 4.5%~4.9%, forming a semi-coherent interface. On the other hand, the spinel interface shell and the α-Fe matrix achieve topological matching through the orientation relationship of the spinel {111} plane parallel to the α-Fe {110} plane. The period of its oxygen sublattice subcells is similar to the period of α-Fe along the

[110] direction, with a mismatch of approximately 3%~6%, also forming a semi-coherent interface. Thus, the original non-coherent α-Al2O3 and α-Fe interface with a mismatch greater than 15% is decomposed into two semi-coherent interfaces with a mismatch of no more than 6%. The interface thus possesses mechanical compliance, enabling coordinated deformation under cyclic loading to suppress interface dissociation. Furthermore, the good bonding between the spinel interface shell and the matrix suppresses the shear dissolution of oxides in the slip band, thereby transforming the alumina dispersed particles from weak chains of fatigue crack initiation into a well-bonded reinforcement resistant to cyclic dissociation.

[0063] The mechanism of action of magnesium stabilizers is manifested in the synergistic solution of multiple deep-seated technical challenges by a single component. Firstly, pure-phase holmite (FeAl2O4) is unstable at high temperatures. Without stabilizers, the resulting holmite interfacial shell decomposes into alumina and iron oxides during sintering at 1200 °C, causing the interface to revert to an incoherent state, as shown in Comparative Example 3. However, magnesium substitution of tetrahedral iron in holmite significantly inhibits its high-temperature decomposition via a magnetite-type inverse spinel mechanism, allowing the spinel interfacial shell to maintain a semi-coherent interface even after sintering at 1200 °C. Secondly, the introduction of magnesium lowers the interfacial energy between the interfacial shell and the matrix, which is beneficial for the continuous coating and stable existence of the interfacial shell on the α-Al2O3 core surface. Thirdly, magnesium also assists in grain refinement and densification during sintering. These multiple effects are simultaneously achieved by the same trace component, magnesium, demonstrating the synergistic effect of a single added component solving multiple deep-seated technical challenges.

[0064] The controlled reduction window setting is the key process mechanism for obtaining and preserving the aforementioned semi-coherent interface shell in situ. The relationship between the relative reduction rate and temperature is as follows: Figure 3As shown in Comparative Example 5, the hydrogen reduction of bulk iron oxides is thermodynamically highly favorable and can be completed rapidly at 400–600 °C. However, the reduction of the magnesium-stabilized iron-aluminum spinel interface shell requires high activation energy and is strongly temperature-dependent. Within the lower temperature range of 480–600 °C, its reduction kinetics are significantly frozen, with the relative reduction rate being only a fraction to a tenth of that at higher temperatures. Therefore, by terminating the hydrogen reduction within a controlled reduction window that both reduces the bulk iron oxides to α-Fe and freezes the reduction kinetics of the spinel interface shell, a continuous α-Fe matrix phase can be obtained while retaining a continuously coated semi-coherent spinel interface shell. When the reduction temperature exceeds this window, the interface shell reduction is no longer frozen and decomposes, resulting in the loss of the beneficial effects. In the chromium-containing embodiments, chromium partially replaces aluminum in the interface shell to form a (Fe,Mg)(Al,Cr)2O4 solid solution interface layer. On the one hand, this finely adjusts the lattice mismatch between the interface shell and α-Fe to further improve the interface matching. On the other hand, chromium dissolves in the α-Fe matrix to enhance its high-temperature resistance to grain boundary oxidation and promote sintering densification, thereby further improving the high-temperature bending fatigue limit based on the aforementioned core-shell structure.

[0065] From the perspective of the interaction between dislocations and second-phase particles, nanoscale core-shell oxide dispersed particles pin dislocations in the α-Fe matrix and hinder their movement. When dislocations bypass particles, they leave dislocation loops according to the Orovan mechanism, thereby improving the strength of the matrix and maintaining the microstructural stability at high temperatures. Under monotonic loading, pure-phase α-Al₂O₃ dispersed particles can exert the above-mentioned pinning effect; however, under high-temperature cyclic loading, the repeated interaction between dislocations and particles accumulates damage at the particle-matrix interface. For incoherent interfaces, the interfacial bonding is weak and the interfacial energy is high. Damage accumulation easily leads to interfacial dissociation and the initiation of fatigue microcracks. At the same time, some oxide particles are sheared or even dissolved in the reciprocating slip bands, resulting in the local loss of matrix reinforcement and cyclic softening. The magnesium-stabilized semi-coherent spinel interface shell of the present invention reduces interface mismatch and interface energy, and improves interface bonding strength, so that the interface can coordinate the local strain caused by dislocation under cyclic loading and suppress interface dissociation. Furthermore, due to the good bonding between the interface shell and the matrix, it suppresses the shear dissolution of oxides in the slip band, thereby simultaneously alleviating both fatigue crack initiation and cyclic softening damage mechanisms.

[0066] During long-term high-temperature service, the thermal stability of core-shell oxide dispersions is crucial for maintaining high-temperature fatigue resistance. The α-Al₂O₃ core itself possesses high thermal stability, while the magnesium-stabilized iron-aluminum spinel interface shell, after magnesium substitution, retains its spinel structure without decomposition at a sintering temperature of 1200 °C. Together, these factors ensure the dimensional and interfacial stability of the core-shell oxide dispersions during sintering and high-temperature service, preventing coarsening of the dispersions and degradation of the interface. In contrast, the solid solution elements relied upon for alloying strengthening diffuse more rapidly at high temperatures, easily leading to grain growth and coarsening of precipitates, resulting in softening. The core-shell oxide dispersions of this invention, due to their high thermal stability, do not suffer from these problems, thus continuously maintaining pinning effects on dislocations and grain boundaries at high temperatures.

[0067] In the case of powder metallurgy gears, the core-shell architecture of this invention is also coupled with the evolution of residual porosity. The magnesium introduced by the magnesium-stabilized spinel interface shell assists in grain refinement and promotes densification during sintering, helping to reduce the size and aggregation of residual porosity in the surface and subsurface layers of the sintered body, thereby reducing stress concentration at the pores. In the chromium-containing embodiment, chromium dissolved in the α-Fe matrix further promotes sintering densification and enhances the high-temperature resistance to grain boundary oxidation of the matrix. The suppression of interface dissociation at the oxide interface and the reduction of stress concentration at residual porosity work together to improve the bending fatigue limit of the powder metallurgy gear under high-temperature heavy-load cyclic conditions. It is precisely due to the synergy of the aforementioned interface engineering and porosity control that the powder metallurgy gear manufactured by this invention achieves a significantly improved high-temperature bending fatigue limit while maintaining high-temperature bending strength.

[0068] It is worth noting that the technical solution of this invention overcomes a technical bias in the field. In the field of oxide dispersion strengthening, those skilled in the art generally believe that the dispersed phase should be an oxide with the highest thermodynamic stability, the greatest hardness, and the strongest chemical inertness to ensure that it is not reduced or dissolved by the matrix at high temperatures and to pin dislocations to the maximum extent. Accordingly, pure phase α-Al2O3 is generally regarded as the preferred dispersed phase due to its high hardness and high thermal stability. However, this invention does not simply pursue high hardness and high stability of the dispersed phase, but intentionally introduces a magnesium-stabilized iron-aluminum spinel interface shell with lower hardness and lower thermodynamic stability than α-Al2O3 on the surface of the α-Al2O3 core. Unexpectedly, this relatively softer and relatively less stable interface shell not only does not weaken the dispersion strengthening effect, but also significantly improves the bending fatigue limit under high-temperature cyclic loading by constructing a semi-coherent interface, improving the interfacial bonding strength, and suppressing interfacial dissociation and shear dissolution. This technical effect, which contradicts conventional design rules, reflects the non-obviousness of this invention.

[0069] In summary, the high-temperature bending fatigue limit of the various embodiments of the present invention is significantly improved compared to the existing wet chemical co-precipitation route, and the bending strength retention rate of powder metallurgy gears at 150 °C is not less than 87% of their room temperature value. Furthermore, the stress-life curve of the 150 °C bending fatigue is generally shifted upwards compared to Comparative Example 1, which uses pure alumina dispersed iron powder. Figure 3 As shown, this provides a high-temperature resistant core-shell alumina dispersion-reinforced iron-based composite powder and its co-precipitation-controlled reduction preparation scheme for high-temperature heavy-duty powder metallurgy gears.

[0070] The above description is only a preferred embodiment of the present invention and is not intended to limit the present invention. Any modifications, equivalent substitutions and improvements made within the spirit and principles of the present invention should be included within the protection scope of the present invention.

Claims

1. A nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears, characterized in that, The composite powder comprises an α-Fe continuous matrix phase and core-shell oxide particles dispersed within the α-Fe continuous matrix phase; the core-shell oxide particles comprise an α-Al₂O₃ core and a magnesium-stabilized iron-aluminum spinel interface layer as the shell, wherein the magnesium-stabilized iron-aluminum spinel interface layer is composed of Fe. (1-x) Mg x Al2O4 with x of 0.1~0.5 is continuously coated on the surface of the α-Al2O3 core by in-situ reaction generated during thermal decomposition. The magnesium-stabilized iron-aluminum spinel interface layer and the α-Fe continuous matrix phase form a semi-coherent interface with a lattice mismatch of no more than 6%. Based on the total mass of the composite powder, the mass fraction of α-Fe is not less than 98%, the mass fraction of α-Al2O3 is 0.45%~1.79%, and the magnesium content is 0.01%~0.4% based on the magnesium-iron molar ratio. The particle size of the core-shell oxide dispersed particles is 10~30 nm.

2. The nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears according to claim 1, characterized in that, The thickness of the magnesium-stabilized iron-aluminum spinel interface layer is 1~3 nm, and it satisfies the orientation relationship between the spinel {111} plane and the α-Fe {110} plane of the α-Fe continuous matrix phase.

3. The nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears according to claim 2, characterized in that, The Fe (1-x) Mg x In Al2O4, x is 0.25~0.35, and the magnesium-stabilized iron-aluminum spinel interface layer maintains the spinel structure without decomposition at a sintering temperature of 1200 °C through a magnetite-type inverse spinel mechanism.

4. The nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears according to claim 3, characterized in that, The magnesium-stabilized iron-aluminum spinel interface layer accounts for 5% to 20% of the total volume fraction of α-Al2O3 in the core-shell oxide dispersed particles.

5. The nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears according to claim 4, characterized in that, The molar ratio of aluminum to iron in the composite powder is 0.5% to 2.0%, and the average spacing between the core-shell oxide dispersed particles in the α-Fe continuous matrix phase is 30 to 50 nm.

6. The nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears according to claim 5, characterized in that, The composite powder also contains chromium, which partially replaces the aluminum in the magnesium-stabilized iron-aluminum spinel interface layer to form a (Fe,Mg)(Al,Cr)2O4 solid solution interface layer and is simultaneously dissolved in the α-Fe continuous matrix phase. The chromium content is 0.1% to 1.0% based on the chromium-iron molar ratio.

7. The method for preparing nano-alumina dispersion-reinforced iron-based composite powder for powder metallurgy gears as described in claim 6, characterized in that, Includes the following steps: Step 1: Dissolve ferric nitrate, aluminum nitrate, and magnesium salt in deionized water to prepare a mixed salt solution. Co-precipitate with ammonia water to simultaneously precipitate ferric hydroxide, aluminum hydroxide, and magnesium hydroxide to obtain a mixed iron-aluminum-magnesium hydroxide precursor. The pH at the end of the coprecipitation is controlled at 9.5~10.

0. Step 2: After washing the precursor, thermally decompose it at a temperature that allows iron oxide and aluminum oxide to react at the interface to form the magnesium-stabilized iron-aluminum spinel interface layer, to obtain an intermediate product with a core-shell oxide as the core and the magnesium-stabilized iron-aluminum spinel interface layer as the shell, dispersed in an iron oxide matrix. Step 3: Reduce the intermediate product with hydrogen under a controlled reduction window that reduces the iron oxide to α-Fe while retaining the magnesium-stabilized iron-aluminum spinel interface layer to obtain the iron-based composite powder.

8. The preparation method according to claim 7, characterized in that, The controlled reduction window in step three is as follows: reduction temperature 480~600 °C, reducing gas dew point -40~-20 °C, and reduction time 30~60 min.

9. The preparation method according to claim 8, characterized in that, The thermal decomposition temperature in step two is 800~900°C, and the precursor is sent into the thermal decomposition furnace after being ultrasonically sprayed for thermal decomposition.

10. A powder metallurgy gear, characterized in that, The powder metallurgical gear is prepared by pressing and sintering the iron-based composite powder according to any one of claims 1 to 6, and the flexural strength retention rate of the powder metallurgy gear at 150 °C is not less than 87% of its room temperature value.

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