Casting al-si alloy and method for producing the same

Abstract

The application provides a cast aluminum-silicon alloy and a preparation method thereof. The cast aluminum-silicon alloy comprises 7.6-10% of Si, 1.0-2.0% of Cu, 0.1-0.3% of Mn, 0.2-0.4% of Mg, 0.052-0.18% of Zr, 0.075-0.165% of V, 0.1-0.15% of Ti, 0.05-0.2% of Ni, 0.01-0.05% of Sn, 0.01-0.05% of Fe, 0.01-0.05% of Sr, and the balance of Al and inevitable impurity elements, by weight percentage; wherein the total amount of the impurity elements is less than or equal to 0.12%. The application refines the grain size through multi-element micro-alloying composition, and the obtained alloy has high room-temperature strength, toughness and high-temperature strength, and is especially suitable for high-end fields such as aerospace, weapon equipment and high-speed railway.

Description

Technical Field

[0001] This invention relates to the field of aluminum alloy technology, and more specifically, to a cast aluminum-silicon alloy and its preparation method. Background Technology

[0002] Cast Al-Si aluminum alloys, with their comprehensive advantages such as low density, good fluidity, low coefficient of thermal expansion, high specific strength, wear resistance, and corrosion resistance, have been widely used in automotive parts engineering. This has significantly reduced vehicle weight, improved fuel economy, and effectively reduced harmful gas emissions. Simultaneously, these alloys exhibit excellent high-temperature resistance, oxidation resistance, and high specific strength, making them widely used in critical heat-resistant components operating under high-temperature conditions in the automotive, aerospace, and marine industries. They have become one of the mainstream materials for next-generation lightweight heat-resistant components.

[0003] However, with the continuous increase in the requirements for component service temperature and durability in the aforementioned fields, the performance margins of traditional cast heat-resistant Al-Si alloys in terms of high-temperature strength, thermal fatigue resistance, and microstructural stability are approaching their limits, making it difficult to meet the development needs of new high-power-density power systems and components operating under extreme conditions. Existing commercial cast aluminum alloys exhibit significant degradation in mechanical properties under high-temperature conditions and are not yet fully capable of withstanding increasingly demanding service environments. To ensure the reliable operation of core components such as engine cylinder heads and turbine housings under long-term exposure to high thermal loads, high stress, and alternating thermal fatigue loads, it is urgent to develop new cast aluminum alloy materials that combine high strength and toughness with excellent high-temperature stability.

[0004] To improve the high-temperature mechanical behavior of Al-Si alloys, researchers have conducted extensive research on multi-element microalloying. In the Al-Si-Cu-Mg alloy system, traditional age-hardening phases tend to coarsen rapidly at high temperatures, leading to a sharp degradation of alloy properties. Cu and Mg are the most important strengthening elements in this system, and their relative contents play a decisive role in the type and distribution characteristics of precipitated phases: when the Cu / Mg ratio is low, the Mg2Si phase dominates, resulting in insufficient high-temperature stability; when the ratio is high, excessive θ-Al2Cu phase precipitation will impair room-temperature plasticity. Therefore, achieving synergistic optimization between high-temperature strength and room-temperature toughness has been a long-standing and unresolved technical challenge in this field.

[0005] Microalloying of transition metals Zr, V, and Ti has been proven to significantly improve the high-temperature performance of Al-Si alloys. These elements possess low diffusion rates and solid solubility in the aluminum matrix, and are considered among the most effective additives for enhancing the alloy's resistance to softening at high temperatures. The Zr, V, and Ti intermetallic compound precipitates they form can exert a significant strengthening effect at high temperatures. However, current techniques struggle to induce the formation of a single, homogeneous Al3(Zr,V,Ti) nanoprecipitate coherent with the aluminum matrix, thus its high-temperature strengthening potential remains largely untapped.

[0006] Ni and Sn are also used to improve the heat resistance of aluminum alloys. Ni can form highly thermally stable intermetallic compounds (such as Al3Ni and Al3CuNi), while Sn can regulate vacancy concentration and promote the uniform precipitation of fine dispersed phases. However, existing methods mostly use individual additions, failing to fully explore the synergistic strengthening potential of the two, thus limiting the effective exertion of the gain effect of this composite microalloying system.

[0007] In summary, the development of a novel cast Al-Si alloy that combines excellent room-temperature mechanical properties with high-temperature heat resistance is urgently needed in the market and has significant engineering application value. Summary of the Invention

[0008] The main objective of this invention is to provide a cast aluminum-silicon alloy and its preparation method, so as to solve the problem that the mechanical properties and high-temperature heat resistance properties of cast aluminum-silicon alloys in the prior art cannot be simultaneously achieved.

[0009] To achieve the above objectives, according to one aspect of the present invention, a cast aluminum-silicon alloy is provided, comprising, by weight percentage, 7.6-10% Si, 1.0-2.0% Cu, 0.1-0.3% Mn, 0.2-0.4% Mg, 0.052-0.18% Zr, 0.075-0.165% V, 0.1-0.15% Ti, 0.05-0.2% Ni, 0.01-0.05% Sn, 0.01-0.05% Fe, 0.01-0.05% Sr, with the balance being Al and unavoidable impurity elements; wherein the total amount of impurity elements is ≤0.12%.

[0010] Furthermore, the weight ratio of Cu to Mg is (4~6):1; and / or the weight ratio of Zr, V to Ti is (0.4~1.8):(0.6~1.5):1; and / or the weight ratio of Ni to Sn is (4~10):1; and / or the ratio of the sum of the weights of Zr, V and Ti to the sum of the weights of Ni and Sn is (1.72~5):1.

[0011] Furthermore, by weight percentage, the cast aluminum-silicon alloy comprises 8-9.5% Si, 1.2-1.8% Cu, 0.13-0.25% Mn, 0.25-0.35% Mg, 0.08-0.16% Zr, 0.078-0.15% V, 0.1-0.15% Ti, 0.08-0.18% Ni, 0.02-0.05% Sn, 0.01-0.05% Fe, 0.01-0.04% Sr, with the balance being Al and unavoidable impurity elements.

[0012] According to another aspect of the present invention, a method for preparing the above-mentioned cast aluminum-silicon alloy is provided, comprising the following steps: Step S1, batching, melting and refining are carried out sequentially according to the composition ratio of the cast aluminum-silicon alloy to obtain a cast aluminum-silicon alloy melt; Step S2, the sand mold raw material is 3D printed to obtain a sand mold; a casting model is set for areas with large wall thickness differences and multiple hot spots, and casting is carried out by cast iron or cast aluminum to obtain chills; the chills are sequentially coated with sand and first baked to obtain chills after first baking; a chilling coating is sprayed onto the chills after first baking and the sand mold, and then a second baking is carried out to obtain chills after second baking and baked sand mold; the chills after second baking and baked sand mold are assembled to obtain a combined mold; Step S3, the cast aluminum-silicon alloy melt is subjected to differential pressure casting in the combined mold to obtain a casting blank; Step S4, the surface of the casting blank is regularized to obtain a casting; the casting is sequentially subjected to solution treatment, water quenching and aging treatment to obtain a cast aluminum-silicon alloy.

[0013] Further, in step S2, the sand mold raw materials include silica sand, resin, and curing agent; wherein the average particle size of the silica sand is 90-150 mesh; the amount of resin added is 10-30 mL / Kg; the amount of curing agent added is 4-10 mL / Kg; the resin includes one or more of furan resin, phenolic resin, and urethane resin; the curing agent includes one or more of acid curing agents, gas curing agents, and peroxide curing agents; and / or, the 3D printing thickness is 0.22-0.4 mm, the sand laying speed is 8-15 s / layer, and the gas emission is 8-12 mL / g; and / or, the first baking temperature is 300-400℃, and the time is 1-3 h; and / or, the second baking temperature is 150-200℃, and the time is 1-3 h; and / or, the chilling coating includes one or more of graphite powder coating, zircon powder coating, and chromite powder coating; and / or, the spraying amount of the chilling coating is 1.0-1.8 g / m 2 .

[0014] Further, in step S3, the temperature of differential pressure casting is 675~695℃; and / or, the synchronous pressure of differential pressure casting is 0.5~0.65MPa; and / or, step S3 also includes a step of ultrasonically treating the cast aluminum-silicon alloy melt simultaneously with differential pressure casting; the power of ultrasonic treatment is 1~2.5KW; and / or, step S3 includes: placing the cast aluminum-silicon alloy melt in the holding furnace of the differential pressure casting machine, connecting the combined mold and the cast aluminum-silicon alloy melt through a riser pipe, and compressing air through a pressure tank to form pressure on the surface of the cast aluminum-silicon alloy melt, so that the cast aluminum-silicon alloy melt sequentially undergoes rising, filling, pressurizing, holding, and depressurizing along the riser pipe from bottom to top.

[0015] Further, in step S3, the rate of liquid influx is 1.1~1.5 kPa / s, the rate of filling is 0.8~1.1 kPa / s; and / or, the pressurization and holding stage includes a pressurization and shelling stage and a holding and crystallization stage. The pressurization rate of the pressurization and shelling stage is 1.0~1.2 kPa / s, the pressurization value is 12~18 kPa, and the time is 8~15 s; the pressurization value of the holding and crystallization stage is 0.4~0.5 MPa, and the holding time is 300~600 s.

[0016] Further, in step S4, the solution treatment includes a first solution treatment, a second solution treatment, and a third solution treatment performed sequentially; the temperature of the first solution treatment is 360~380℃, and the holding time is 8~14h; the temperature of the second solution treatment is 490~510℃, and the holding time is 3~5h; the temperature of the third solution treatment is 515~535℃, and the holding time is 8~12h; and / or, the water quenching temperature is 60~80℃, and the quenching transfer time is 5~10s; and / or, the aging treatment temperature is 160~175℃, and the holding time is 10~14h.

[0017] Further, step S1 includes: step S11, melting and holding the aluminum ingot at the first holding temperature, and sequentially adding AlSi master alloy, AlCu master alloy, AlNi master alloy, AlMn master alloy, AlV master alloy and Sn ingot at the first holding temperature to obtain a first intermediate melt; step S12, holding the first intermediate melt at the second holding temperature, and sequentially adding AlZr master alloy and Mg ingot at the second holding temperature to melt them, and then adding AlTi master alloy to mix them to obtain a second intermediate melt; passing argon gas through the second intermediate melt for refining, adding a refining agent to obtain a third intermediate melt; adding AlSr master alloy to the third intermediate melt for modification treatment to obtain a cast aluminum-silicon alloy melt.

[0018] Further, in step S11, the weight percentage of Si in the AlSi master alloy is 20-30%; and / or, the weight percentage of Cu in the AlCu master alloy is 30-50%; and / or, the weight percentage of Mn in the AlMn master alloy is 10-20%; and / or, the weight percentage of Zr in the AlZr master alloy is 5-10%; and / or, the weight percentage of V in the AlV master alloy is 5-10%; and / or, the weight percentage of Ti in the AlTi master alloy is 10-15%; and / or, the weight percentage of Ni in the AlNi master alloy is 10-20%; and / or, the melting temperature is 730-750℃, and the time is 0.5-1h; and / or, the first holding temperature is 720-740℃; and / or, the second holding temperature is 700℃. ~720℃; and / or, step S11 further includes the following steps: before melting, aluminum ingots, AlSi master alloy, AlCu master alloy, AlMn master alloy, AlZr master alloy, AlV master alloy, AlTi master alloy, AlNi master alloy, Sn ingots and Mg ingots are respectively subjected to a third baking; the temperature of the third baking is 150~250℃, and the time is 1~3h; and / or, in step S12, the refining temperature is 730~740℃, the time is 20~30min, and the rotation speed is 400~600r / min; and / or, the refining agent includes one or more of AlTiB, AlTiC and TCB, and the weight ratio of the refining agent to the first intermediate melt is 0.3~0.8%; and / or, the temperature of the modification treatment is 720~740℃, and the time is 20min~40min.

[0019] By employing the technical solution of this invention, and through the use of multi-element microalloying components such as Si, Cu, Mn, Mg, Zr, V, Ti, Ni, Sn, Fe, and Sr, the grain size is significantly refined. The resulting cast aluminum-silicon alloy exhibits comprehensive improvements in room temperature tensile strength, elongation, and high temperature tensile strength and elongation, thus possessing high room temperature strength, high toughness, and excellent high temperature strength. It has significant industrial practical value and broad application prospects, and is particularly suitable for high-end fields such as aerospace, weaponry, and high-speed rail. Attached Figure Description

[0020] The accompanying drawings, which form part of this application, are used to provide a further understanding of the invention. The illustrative embodiments of the invention and their descriptions are used to explain the invention and do not constitute an undue limitation of the invention. In the drawings:

[0021] Figure 1 The electron backscatter diffraction (EBSD) microstructure of the alloy of Example 1 of the present invention in the T6 state is shown;

[0022] Figure 2The image shows the metallographic structure (OM) of the alloy of Example 1 of the present invention in the as-cast state;

[0023] Figure 3 A transmission electron microscope (TEM) image of the alloy of Example 1 of the present invention in the T6 state is shown;

[0024] Figure 4 The transmission electron microscopy diffraction pattern of the alloy of Example 1 of the present invention in the T6 state is shown. Detailed Implementation

[0025] It should be noted that, unless otherwise specified, the embodiments and features described in this application can be combined with each other. The present invention will now be described in detail with reference to the accompanying drawings and embodiments.

[0026] As described in the background section of this invention, existing technologies suffer from the problem of being unable to simultaneously achieve both high-temperature heat resistance and mechanical properties in cast aluminum-silicon alloys. To address this issue, in a typical embodiment of this invention, a cast aluminum-silicon alloy is provided, comprising, by weight percentage: 7.6-10% Si, 1.0-2.0% Cu, 0.1-0.3% Mn, 0.2-0.4% Mg, 0.052-0.18% Zr, 0.075-0.165% V, 0.1-0.15% Ti, 0.05-0.2% Ni, 0.01-0.05% Sn, 0.01-0.05% Fe, 0.01-0.05% Sr, with the balance being Al and unavoidable impurity elements; wherein the total amount of impurity elements is ≤0.12%. The content of a single unavoidable impurity element is ≤0.05%, and the unavoidable impurity elements include one or more of C, P and S.

[0027] Si is the most essential basic alloying element in Al-Si cast aluminum alloys. Si plays multiple roles in aluminum alloys. First, Si forms a eutectic reaction with Al, precipitating a large number of fine eutectic silicon phases during solidification. These eutectic silicon phases are distributed in the α-Al matrix in a fibrous or granular form, acting as a hard, wear-resistant phase to improve the alloy's wear resistance and contributing to strengthening by hindering dislocation movement. Second, Si atoms dissolved in the α-Al matrix cause lattice distortion, providing solid solution strengthening. More importantly, Si is a key component in the formation of aging precipitates. It combines with Mg to form the Mg2Si strengthening phase and forms the Q phase with Cu and Mg. These nanoscale precipitates are uniformly dispersed during aging, working synergistically with the θ′ phase to achieve significant precipitation strengthening. Furthermore, Si effectively reduces the alloy's coefficient of thermal expansion, improving the dimensional stability and thermal fatigue resistance of castings.

[0028] When the Si content is within the above range, the alloy exhibits excellent casting fluidity, significantly improving the filling capacity of complex thin-walled castings and effectively reducing the tendency for hot cracking. Simultaneously, the moderate amount of eutectic silicon significantly improves wear resistance, maintains a low coefficient of thermal expansion, and provides good dimensional stability. More importantly, the ample Si supply ensures the precipitation of Mg2Si and Q phases, fully leveraging the precipitation strengthening effect and ensuring the alloy's high-temperature strength fully meets design requirements. Furthermore, the above-mentioned Si content ensures a moderate amount of eutectic silicon, which, after Sr modification treatment, exhibits a fine fibrous distribution. This provides sufficient strengthening and wear resistance to the matrix without severely disrupting its continuity, thus guaranteeing the required room-temperature elongation. Additionally, this Si content range ensures a reasonable ratio with Cu and Mg, providing sufficient precipitation driving force for Mg2Si and Q phases, forming a multiple precipitation strengthening effect with the θ′ phase, and fully releasing the alloy's precipitation strengthening potential. In addition, the synergistic effect of Si and Mn can transform the needle-like phase formed by Fe impurities into blocky or Chinese character-shaped phases, effectively eliminating the cutting damage of needle-like phases to the matrix, and enabling the alloy to maintain excellent plasticity and toughness within the allowable Fe content range.

[0029] Cu, as one of the most important strengthening elements in the Al-Si alloy of this invention, has its content strictly controlled within the aforementioned range. This content range was determined through extensive experimental optimization, taking into full account precipitation strengthening effects, thermal stability of precipitated phases, grain boundary segregation tendency, casting performance, and synergistic effects with other key elements such as Si and Mg. Copper exhibits multiple strengthening mechanisms in the aluminum matrix. First, during solution aging, Cu precipitates a high-density, dispersed metastable θ′ (Al2Cu) phase. This phase is disc-shaped, coherent or semi-coherent with the aluminum matrix, and uniformly disperses during aging, producing a strong precipitation strengthening effect, which is the main source of the alloy's high room temperature strength. Second, Cu is an indispensable component of the Q phase. The addition of Cu causes the precipitated phase in the alloy to change from a Mg2Si-dominated type to a synergistic strengthening type of Q and θ′ phases. The Q phase has a body-centered cubic structure, a high coarsening activation energy, and exhibits significantly better thermal stability than the Mg2Si and θ′ phases in the mid-temperature range of 200–350 °C, allowing it to continuously exert a strengthening effect under high-temperature, long-term service conditions. In addition, Cu atoms dissolved in the α-Al matrix cause lattice distortion, providing additional solid solution strengthening.

[0030] When the Cu content is higher than 1.0 wt.%, the precipitation driving force of the θ′ phase and Q phase is sufficient, the volume fraction of the precipitated phase is high, and the precipitation strengthening effect can be fully released, making it easy for the room temperature and high temperature tensile strength of the alloy to meet the design requirements. This invention controls the Cu content within the above range, ensuring sufficient precipitation driving force and appropriate precipitation quantity for the θ′ and Q phases, while controlling the microsegregation of Cu to an acceptable level, avoiding the formation of continuous brittle network θ phase at grain boundaries, thus maximizing the precipitation strengthening effect. This invention precisely controls the Cu content within the above range, ensuring that the precipitation strengthening of Cu and its synergistic effect with the Q phase play a core supporting role in achieving these performance indicators. When the Cu content is within the above range, the θ′ phase and Q phase have sufficient precipitation driving force and appropriate precipitation quantity, maximizing the precipitation strengthening effect; simultaneously, the microsegregation of Cu is controlled to an acceptable level, avoiding the formation of continuous brittle network θ phase at grain boundaries, thus maintaining a high level of room temperature elongation and hot crack resistance.

[0031] Mn is an important auxiliary alloying element in the Al-Si alloy of this invention and is the most effective element for neutralizing the harmful effects of Fe impurities. In cast Al-Si alloys, Fe is usually introduced unavoidably as an impurity element from raw materials and the smelting process. The addition of Mn can change the crystallization behavior of Fe-containing phases. Mn preferentially participates in the nucleation and growth of Fe-rich phases, transforming acicular phases into Chinese character-shaped, skeletal, or blocky phases. This phase has a rounded morphology and uniform distribution, significantly reducing its cutting effect on the matrix, thereby effectively eliminating the harm of Fe impurities. Mn provides supplementary solid solution strengthening. Mn has a certain solid solubility in the aluminum matrix. Mn atoms dissolved in α-Al cause lattice distortion, resulting in a solid solution strengthening effect. The atomic radius of Mn differs from that of Al, and the distorted stress field it causes can effectively hinder dislocation movement. This, combined with the precipitation strengthening effect provided by elements such as Cu and Mg, further improves the room temperature and high temperature strength of the alloy. At the same time, Mn has a low diffusion rate in aluminum, and its solid solution strengthening effect decays slowly at high temperatures, contributing to the high-temperature performance of the alloy. Furthermore, an appropriate amount of Mn can interact with Si to change the growth mode and interface characteristics of eutectic silicon, making the eutectic silicon distribution more uniform and its morphology more refined. This has a synergistic promoting effect on the fine fibrous distribution of eutectic silicon in the alloy after Sr modification treatment, which is beneficial to improving the plasticity and toughness of the alloy.

[0032] When the Mn content is within the above-mentioned range, it has a significant effect on modifying the harmful acicular Fe phase, transforming the acicular β phase into a less harmful blocky or Chinese character-shaped phase, effectively eliminating the severing damage to the matrix and thus ensuring the alloy's plasticity. Simultaneously, Mn provides sufficient solid solution strengthening, and the resulting lattice distortion effectively hinders dislocation movement, significantly improving high-temperature strength. Within this range, the problem of excessive Mn forming coarse Mn-containing intermetallic compound phases is avoided, preventing these coarse phases from becoming stress concentration sources and crack initiation points, ensuring the alloy's high plasticity and long fatigue life. Furthermore, an appropriate amount of Mn does not alter the alloy's solidification path, controls the width of the mushy region, and is beneficial for the effective control of shrinkage defects during differential pressure casting, improving the density of the casting.

[0033] Mg is a core component of age-hardening precipitation in Al-Si alloys. Mg can combine alone with Si to form the Mg2Si strengthening phase, or it can combine with Cu and Si to form the Q phase. The Q phase has a body-centered cubic structure, and its coarsening activation energy is significantly higher than that of the θ′ and β″ phases. It exhibits excellent thermal stability in the mid-temperature range of 200–350 °C, making it the most effective strengthening phase in this temperature range. Simultaneously, Mg has a high solid solubility in the aluminum matrix, far exceeding the room-temperature solid solubility of Cu and Si. The Mg atoms dissolved in the α-Al matrix have a significant atomic size difference with the aluminum matrix, resulting in strong lattice distortion and forming a high-density elastic stress field. This effectively hinders dislocation movement and contributes significantly to solid solution strengthening. This solid solution strengthening effect plays a crucial role in ensuring basic strength before aging treatment and improving overall performance after aging. More importantly, the formation of the Q phase consumes some Mg and Si, regulating the ratio of the θ′ and Q phases in the alloy, allowing them to complement each other in terms of quantity density and distribution morphology, jointly constituting the main source of precipitation strengthening in the alloy. Furthermore, an appropriate amount of Mg can alter the growth interface energy of eutectic silicon, to some extent assisting the effect of Sr modification treatment, resulting in higher fibrousness, finer spacing, and more uniform distribution of eutectic silicon, which is beneficial to improving the alloy's plasticity and toughness.

[0034] When the Mg content is within the above range, the precipitation driving force of Mg2Si and Q phases is sufficient, and the precipitation strengthening effect is effectively established. At the same time, Mg has a large solid solubility in the aluminum matrix, producing strong lattice distortion, contributing a significant solid solution strengthening effect, and making the alloy's aging response rapid, with comprehensive mechanical properties meeting service requirements. Within this range, the problem of Mg2Si phase dominance can be avoided, preventing the degradation of strength and heat resistance caused by the high coarsening rate of Mg2Si phase at high temperatures. At the same time, an appropriate amount of Mg does not aggravate grain boundary segregation during solidification, resulting in a low tendency for hot cracking. In addition, a moderate Mg content can avoid the formation of coarse Mg-rich phases, ensuring that Cu and Si elements fully participate in the precipitation of Q phase, maintaining the balanced precipitation of multi-component strengthening phases, improving strengthening efficiency, and ensuring alloy plasticity.

[0035] Zr is a key microalloying and high-temperature stabilizing element in the Al-Si alloy of this invention. Zr has an extremely low diffusion rate and limited solid solubility in the aluminum matrix. During post-solidification aging or high-temperature service, Zr precipitates from the supersaturated solid solution as the Al3Zr phase. This precipitated phase has an L12 ordered structure, is completely coherent with the aluminum matrix, and has a lattice mismatch of less than 1%. The coherent strain field effectively hinders dislocation movement and grain boundary migration, providing significant precipitation strengthening. More importantly, the Al3Zr phase has an extremely high coarsening activation energy, maintaining nanoscale size and good coherence even above 475℃. Its coarsening rate is several orders of magnitude lower than that of the traditional strengthening phases θ′ and Mg2Si, which is crucial for ensuring the performance stability of the alloy under high-temperature, long-term service conditions. Furthermore, the fine and dispersed Al3Zr nanophase effectively pins grain boundaries and subgrain boundaries, inhibiting recrystallization and grain growth, further maintaining the microstructure stability of the alloy at high temperatures.

[0036] When the Zr content is within the above range, the precipitation driving force of Al3(Zr,V,Ti) nanophase is sufficient, and the number density of precipitates per unit volume is high, which can form sufficient pinning force to effectively stabilize grain boundaries and subgrain boundaries. Within this range, the problem of excessive Zr forming primary Al3Zr coarse particles during solidification due to limited solid solubility is avoided, preventing the coarse phase from losing its coherent relationship with the matrix and ensuring that the risk of it acting as a stress concentration source and crack initiation point is eliminated, thereby ensuring the plasticity and fatigue performance of the alloy. At the same time, an appropriate amount of Zr does not change the normal nucleation and growth behavior of Al3(Zr,V,Ti) precipitates, maintaining the balance of the Zr / V / Ti ratio in the multi-component composite phase, ensuring the formation conditions of a single homogeneous phase, and giving full play to the high-temperature strengthening effect.

[0037] Ti, as a key microalloying and nucleation promoting element in the Al-Si alloy of this invention, is the core nucleation element that promotes the fine and uniform precipitation of the multi-component composite nanophase. The introduction of Ti leads to the co-entry of Zr, V, and Ti into the ordered lattice of L12, forming a single homogeneous nanoprecipitate. In the formation of this multi-component composite phase, Ti plays a crucial role in lowering the nucleation energy barrier and promoting the uniform precipitation of a large number of fine nuclei. The atomic size and electronic structure of Ti have good compatibility with the Zr sites in the Al3Zr phase. When Ti partially replaces Zr, the lattice constant of the precipitate is finely adjusted, and the mismatch with the aluminum matrix is ​​further reduced, thereby reducing the interfacial energy required for the nucleation of the new phase, decreasing the critical nucleation radius, and significantly improving the nucleation rate of the precipitate. This mechanism enables the nanophase to precipitate uniformly with a higher number density in a shorter aging time, avoiding the problems of insufficient quantity and uneven distribution of precipitates caused by the difficulty of nucleation when Zr is added alone. Meanwhile, the fine and dispersed precipitate nuclei provide abundant nucleation sites for the subsequent diffusion and aggregation of Zr and V, thus comprehensively optimizing the precipitation kinetics of the multi-component composite phase.

[0038] During melt solidification, Ti and Al can form the Al3Ti phase, which has the same tetragonal lattice structure (D0) as α-Al. 22 The equiaxed α-Al (a type of α-Al) has a lattice constant similar to that of the aluminum matrix, making it a suitable heterogeneous nucleation site for α-Al and significantly refining the as-cast grain structure. This grain refinement, on the one hand, reduces the dendrite arm spacing of α-Al, resulting in a more uniform and dispersed distribution of eutectic silicon and intermetallic compounds, and reducing the segregation of coarse second phases at grain boundaries. On the other hand, the fine equiaxed grain structure improves the feeding ability of castings, reduces casting defects such as shrinkage porosity and hot cracks, and provides a good grain base for obtaining a uniform precipitate structure during subsequent heat treatment. Furthermore, the Ti atoms dissolved in the aluminum matrix can produce localized lattice distortion, increasing the resistance to dislocation movement and providing a certain solid solution strengthening effect. Especially at high temperatures, the low diffusion rate of Ti causes its solid solution strengthening effect to decay slowly, thus contributing to the improvement of the alloy's high-temperature mechanical properties.

[0039] When the Ti content is within the above range, the nucleation driving force and nucleation rate of the nanophase are sufficient, the number density of precipitates is high and the size is small, which can form a fully dispersed nano-reinforcing network. In addition, the grain refinement effect is significant, the as-cast structure is fine, which is conducive to the homogenization control of the subsequent structure and enables the high-temperature strength to meet the design requirements. Within this range, the problem of excessive Ti forming coarse primary Al3Ti phases during solidification is avoided, and the coarse phases are prevented from severely cutting the aluminum matrix due to their sharp edges, eliminating the risk of them acting as stress concentration sources and crack initiation points, thereby ensuring the room temperature and high-temperature plasticity of the alloy. More importantly, the appropriate amount of Ti does not increase the density of the alloy or the tendency for hot cracking, and has a promoting effect on the forming quality of complex thin-walled castings in differential pressure casting.

[0040] Ni is a key heat-strengthening element and a crucial element for ensuring performance in the Al-Si alloy of this invention within the mid-temperature range (200~350℃). Ni has a low diffusion rate and extremely limited solid solubility in the aluminum matrix. During solidification and heat treatment, it reacts with Al and Cu to form intermetallic compounds such as Al3Ni and Al3CuNi. This ensures the aluminum alloy remains stable in the solid state during solution treatment and high-temperature service, without dissolution or significant coarsening. Simultaneously, the aforementioned Ni-rich phase imparts high hardness and good thermal stability to the aluminum alloy, effectively hindering dislocation movement and grain boundary sliding within the mid-temperature range of 200~350℃, providing significant resistance to high-temperature softening. Furthermore, the Ni-rich phase forms a stable coherent or semi-coherent interface with the aluminum matrix, resulting in low interfacial energy, low coarsening driving force, and maintaining its strengthening effect even under long-term high-temperature service conditions.

[0041] Ni and Sn work synergistically to achieve highly refined and homogenized Ni-rich phases through a heterogeneous nucleation mechanism. When Ni and Sn are added together, Sn's "vacancy engineering" (referring to a material modification method that intentionally introduces, regulates, and manages atomic vacancies in crystalline materials through controllable means to directionally optimize their electronic structure, surface chemical properties, and other functional properties) not only promotes the precipitation of all dispersed phases, but more importantly, because Sn has extremely low solid solubility and a low melting point in the aluminum matrix, Sn preferentially accumulates in the matrix to form nanoscale Sn-rich clusters or particles during the early stages of solidification and aging. These Sn-rich nanophases have low-energy interfaces and good lattice matching with the subsequently precipitated Al3Ni and Al3CuNi phases, and can serve as efficient heterogeneous nucleation cores for Ni-rich phases, guiding the uniform and fine precipitation and growth of Ni-rich phases around them. This alters the intrinsic tendency of Ni-rich phases to form coarse lamellar or needle-like morphologies when precipitated alone, significantly increasing the number density of Ni-rich strengthening phases, refining their size to the nanoscale, and greatly improving their distribution uniformity, ultimately significantly enhancing the high-temperature strengthening effect of Ni. Furthermore, the dispersed distribution of Ni-rich intermetallic compound hard particles within the aluminum matrix effectively improves the alloy's wear resistance and reduces the coefficient of friction. Simultaneously, Ni's low coefficient of thermal expansion further reduces the overall coefficient of thermal expansion of the alloy, improving the dimensional stability and thermal fatigue resistance of the aluminum alloy under high-temperature service conditions.

[0042] When the Ni content is within the above range, the number density of Ni-rich intermetallic compounds is sufficient to form an effective strengthening network in the mid-temperature range, exhibiting excellent resistance to high-temperature softening in the 200~350℃ temperature range. Simultaneously, within this range, the problem of excessive Ni forming coarse primary Al3Ni needle-like or lath-like phases during solidification due to limited solubility in the aluminum melt can be avoided. This prevents the coarse phases from severely cutting the aluminum matrix due to their sharp edges, eliminating the risk of them serving as preferential crack initiation and propagation channels, thus ensuring the alloy's room-temperature plasticity and fatigue resistance. Furthermore, an appropriate amount of Ni ensures that the Ni:Sn ratio does not exceed 10, and the number of heterogeneous nucleation cores of Sn is sufficient to refine and disperse the large amount of Ni-rich phases, ensuring that the Ni phase precipitates with a fine morphology. In addition, avoiding excessive Ni consumption by coarse Ni-rich phases controls the alloy density, which is beneficial for weight reduction and has a positive impact on casting fluidity and feeding ability.

[0043] Sn is a key microalloying element and a "vacancy engineering" control element in the Al-Si alloy of this invention. Sn atoms have extremely high binding energy with vacancies in the aluminum matrix. During solution treatment and quenching, Sn atoms can effectively capture and retain a large number of supersaturated vacancies, significantly increasing the vacancy concentration in the quenched matrix. This provides numerous rapid diffusion channels for solute atoms during subsequent aging, reducing the diffusion barriers of elements such as Cu, Mg, and Si, and promoting the uniform and dense nucleation of the GP zone and θ″ phase precursors in the early stages of aging. The increased vacancy concentration also alters the nucleation kinetics of the precipitated phases, allowing nanophases that would normally require a long incubation period to precipitate to form in a shorter time. Sn precipitates uniformly at a higher density, significantly enhancing the alloy's aging response speed and precipitation strengthening effect. Furthermore, appropriate Sn segregation at the solid-liquid interface and grain boundaries reduces interfacial energy, alters solute redistribution behavior at the solidification front, and refines eutectic silicon and intermetallic compounds in the as-cast microstructure to some extent. Simultaneously, Sn segregation at grain boundaries improves the wettability of the second phase with the matrix, alleviates stress concentration, and contributes to improving the alloy's plasticity. In addition, Sn's low melting point promotes atomic diffusion and compositional homogenization during solution treatment.

[0044] When the Sn content is within the above range, the vacancy concentration in the matrix is ​​sufficient, and the "vacancy engineering" effect can be effectively established, significantly promoting the heterogeneous nucleation of the Ni-rich phase, guiding the uniform and fine precipitation of the Ni phase, and resulting in excellent strengthening effect. At the same time, a moderate Sn content keeps the Ni:Sn ratio in the range of 5 to 10, allowing the "catalytic" effect of Sn to be fully utilized, making a significant contribution to mid-temperature strengthening, and meeting the design requirements for high-temperature tensile strength at 300℃. Within this range, it is necessary to avoid the problem that excessive Sn, due to its extremely low solid solubility and low melting point, cannot dissolve in the matrix and will agglomerate in large quantities at grain boundaries and dendrites, preventing the formation of continuous or semi-continuous distribution of the Sn-rich low-melting-point eutectic phase at grain boundaries, eliminating the risk of softening or melting during high-temperature service or heat treatment, which would reduce the grain boundary bonding force, thereby avoiding the "hot brittleness" phenomenon of the alloy, ensuring room temperature elongation and reducing the tendency for hot cracking. At the same time, an appropriate amount of Sn avoids excessive Sn grain boundary agglomeration, which would damage the nucleation and distribution behavior of the Ni-rich phase, ensuring the optimization of the high-temperature strengthening effect.

[0045] Fe plays a distinctly dual role in Al-Si alloys, but overall, in the high-strength and high-toughness alloy system of this invention, Fe is considered an element that must be strictly limited. During melt solidification, Fe undergoes a eutectic reaction with Al and Si to form needle-like or lath-like β-Al5FeSi phases. These phases exhibit sharp needle-like or long lath-like morphologies, with lengths ranging from tens to hundreds of micrometers. The β-Al5FeSi phase has extremely high hardness but extremely poor plasticity, and its interfacial bonding with the aluminum matrix is ​​weak. Furthermore, under stress, severe stress concentration occurs at the tips and edges of the needle-like phases, which easily become channels for preferential crack initiation and propagation. Therefore, with increasing Fe content, the room temperature elongation and fracture toughness of the alloy decrease exponentially. Even if the Fe content increases only from 0.05 wt.% to 0.15 wt.%, the elongation may decrease by 30-50%. Fe has extremely low solid solubility in the aluminum matrix and strongly segregates between dendrites and grain boundaries during solidification. The acicular β-Al5FeSi phase forms at the end of solidification, which severely hinders the feeding channels between dendrites and increases the tendency for shrinkage porosity. At the same time, the brittle acicular phase is distributed in a network at the grain boundaries, which significantly reduces the grain boundary bonding force and increases the susceptibility to hot cracking.

[0046] This invention utilizes the aforementioned behavior mechanism of Fe in Al-Si melt to neutralize the harmful effects of Fe. The addition of Mn can alter the crystallization behavior of the Fe-rich phase: Mn preferentially participates in the nucleation and growth of the Fe-rich phase, transforming the needle-like β-Al5FeSi phase into Chinese character-shaped, skeletal, or blocky α-Al. 15 (Fe,Mn)3Si2 phase. This phase has a rounded morphology and uniform distribution, which greatly reduces its cutting effect on the matrix. At the same time, its interfacial bonding force is better than that of the β phase, which can reduce stress concentration to a certain extent.

[0047] When the Fe content is within the aforementioned range, it balances the economics and process controllability of industrial production. It avoids the significant increase in raw material costs and smelting control difficulties associated with pursuing ultra-low Fe content. Simultaneously, it eliminates the need to compensate for the potential impact of trace Fe deficiency on high-temperature strength by adding other expensive alloying elements, ensuring good economic efficiency. Furthermore, an appropriate Fe content ensures suitable mold release properties and wear resistance. Within this range, the quantity and size of the acicular β-Al5FeSi phase are strictly controlled, ensuring that the modification effect of Mn is sufficient to completely suppress the formation of harmful acicular phases, eliminating the risk of residual harmful acicular phases in the alloy. This avoids the problem of these phases forming a three-dimensional network in the casting body, severely disrupting the matrix and causing a sharp decrease in elongation. At the same time, an Fe content not exceeding 0.05 wt.% avoids the consumption of some Mn and Si elements by the Fe-rich phase, ensuring that the regulatory effect of Mn on harmful phases and the strengthening contribution of Si to the matrix are fully realized.

[0048] In summary, this invention employs a multi-element microalloying composition including Si, Cu, Mn, Mg, Zr, V, Ti, Ni, Sn, Fe, and Sr, resulting in significantly refined grain size. This comprehensively enhances the room-temperature tensile strength, elongation, and high-temperature tensile strength and elongation of the resulting cast aluminum-silicon alloy. Consequently, it possesses high room-temperature strength, high toughness, and excellent high-temperature strength, demonstrating significant industrial practical value and broad application prospects, particularly suitable for high-end fields such as aerospace, weaponry, and high-speed rail.

[0049] In a preferred embodiment, the weight ratio of Cu to Mg is (4~6):1; and / or, the weight ratio of Zr, V to Ti is (0.4~1.8):(0.6~1.5):1; and / or, the weight ratio of Ni to Sn is (4~10):1; and / or, the ratio of the sum of the weights of Zr, V and Ti to the sum of the weights of Ni and Sn is (1.72~5):1.

[0050] Limiting the Cu / Mg mass ratio mentioned above is more conducive to achieving an equilibrium ratio between the Q phase and the θ′ phase in the alloy. The θ′ phase mainly contributes to room temperature strength, while the Q phase ensures high-temperature stability. Within this range, the problem of Mg2Si phase dominance caused by an excessively low Cu / Mg ratio can be further reduced, allowing the alloy to maintain a higher strength level over a wide temperature range from room temperature to 300℃. Simultaneously, it is easier to reduce the problem of excessive θ′ phase precipitation and reduced Q phase proportion caused by an excessively high Cu / Mg ratio, further reducing the risk of impaired high-temperature strengthening effect and the risk of excessive θ′ phase coarsening and transforming into the equilibrium θ phase at high temperatures, thus further degrading high-temperature performance. It also further reduces the problem of a sharp decrease in room temperature elongation and a significant increase in hot cracking tendency caused by the enrichment of coarse equilibrium θ-Al2Cu phase at grain boundaries and dendrites due to excess Cu relative to Mg, thereby further improving the alloy's excellent comprehensive mechanical properties at both room temperature and high temperatures.

[0051] This invention limits the weight ratio of Zr, V, and Ti within the aforementioned range, which is more conducive to controlling the entry of the three elements into the L12 structural lattice and making it easier to form a single homogeneous Al3(Zr,V,Ti) nano-precipitate. This makes it more advantageous for the strengthening efficiency, thermal stability, and precipitation kinetics of the Al3(Zr,V,Ti) nano-phase to be superior to their respective single precipitates. It also further reduces the problem of forming coarse primary phases when a single element is added due to the concentration of a certain element exceeding the solid solution limit, thereby further improving the overall performance of aluminum alloys.

[0052] When the weight ratio of Ni to Sn is within the above range, it is more conducive to the catalytic effect of Sn, thereby further promoting the full manifestation and transformation of the Ni phase, making the contribution of mid-temperature strengthening more significant, and further improving the tensile strength at 300℃.

[0053] Secondly, controlling the Ni:Sn mass ratio at 4~10 is more conducive to increasing the number density of Ni-rich strengthening phases, refining crystal size, and improving the uniformity of Ni and Sn distribution, thereby further enhancing the high-temperature strengthening capability of aluminum alloys.

[0054] To further improve the room temperature tensile strength, elongation, and high temperature tensile strength and elongation of the cast aluminum-silicon alloy, in a preferred embodiment, the cast aluminum-silicon alloy comprises, by weight percentage, 8-9.5% Si, 1.2-1.8% Cu, 0.13-0.25% Mn, 0.25-0.35% Mg, 0.08-0.16% Zr, 0.078-0.15% V, 0.1-0.15% Ti, 0.08-0.18% Ni, 0.02-0.05% Sn, 0.01-0.05% Fe, 0.01-0.04% Sr, with the balance being Al and unavoidable impurity elements.

[0055] In another typical embodiment of the present invention, a method for preparing the above-mentioned cast aluminum-silicon alloy is also provided, comprising the following steps: Step S1, batching, melting and refining are carried out in sequence according to the composition ratio of the cast aluminum-silicon alloy to obtain a cast aluminum-silicon alloy melt; Step S2, the sand mold raw material is 3D printed to obtain a sand mold; a casting model is set for areas with large wall thickness differences and multiple hot spots, and casting is carried out by cast iron or cast aluminum to obtain chills; the chills are sequentially coated with sand and first baked to obtain chills after first baking; a chilling coating is sprayed onto the chills after first baking and the sand mold, and then a second baking is carried out to obtain chills after second baking and baked sand mold; the chills after second baking and baked sand mold are assembled to obtain a combined casting mold; Step S3, the cast aluminum-silicon alloy melt is subjected to differential pressure casting in the combined casting mold to obtain a casting blank; Step S4, the surface of the casting blank is regularized to obtain a casting; the casting is sequentially subjected to solution treatment, water quenching and aging treatment to obtain a cast aluminum-silicon alloy.

[0056] In existing technologies, at the manufacturing process level, traditional sand casting faces problems such as low production efficiency, reliance on experience for process parameters, and insufficient dimensional accuracy of castings. While differential pressure casting can produce high-quality castings with dense microstructure and excellent performance, its integration with sand mold 3D printing technology still requires systematic solutions to the synergistic matching of process steps such as mold structure optimization, chiller placement strategies, and filling and pressurization regimes. Furthermore, current processes still have many shortcomings in melt purification, filling process control, and heat treatment regime design, hindering the effective conversion of the alloy's intrinsic properties into actual castings.

[0057] To further address the aforementioned process problems, this invention provides a preparation method. First, the raw materials are sequentially batched, smelted, and refined to remove gases and inclusions from the melt, resulting in a homogeneous and clean cast aluminum-silicon alloy melt. Then, numerical simulation is used to predict the casting filling process, solidification sequence, and shrinkage porosity distribution. Based on this, the optimal layout of the gating system, risers, and chills is designed, pre-verifying the feasibility of the process scheme in digital space and avoiding the high cost and long cycle of traditional trial-and-error methods. For multi-element alloys containing elements with high segregation tendency such as Zr, V, and Ti, simulation optimization can effectively predict element segregation regions and guide process adjustments to achieve microstructure homogeneity. Finally, 3D printing technology is used to construct sand mold structures with complex geometries by layering the raw materials.

[0058] Then, a casting model is set up for areas with large wall thickness differences and multiple hot spots, and the chills are obtained by casting with cast iron or cast aluminum.

[0059] When a casting has thin-walled stiffeners (e.g., 3-5 mm thick) directly connected to a thick cross-section (e.g., a wall seat or flange 15-25 mm thick), thermal stress concentration is easily generated at the interface due to the significant difference in cooling rates; this area is known as the region with a large wall thickness difference. A complex geometric region on the casting outline with poor local heat dissipation conditions, a significantly longer solidification time than surrounding areas, and two or more overlapping hot spots is called a multi-hot spot region. Historical data after X-ray inspection or cutting reveals a high likelihood of shrinkage porosity or shrinkage defects. Installing chills in regions with large wall thickness differences and multiple hot spots can accelerate local heat dissipation, eliminate hot spots, and achieve sequential solidification, thereby reducing casting defects.

[0060] Sand coating is applied to the chill. Sand coating involves covering the surface of the chill with a thin layer of sand. The purpose is to prevent cold shuts or chilling cracks from forming on the casting surface due to excessive cooling when the chill is in direct contact with molten aluminum. The sand coating provides moderate thermal resistance, thus moderating the chilling effect. This ensures sufficient cooling rate to refine the microstructure while avoiding casting stress caused by extreme temperature gradients.

[0061] After the chills undergo sand coating, a sand layer with specific thermal resistance is formed on their surface. The first baking process removes moisture and volatiles from the chills and sand layer. A chilling coating is sprayed on, and after a second baking, the chill surface acquires a coating with good thermal conductivity and surface integrity, while the sand mold remains dry. The second-baked chills are then assembled with the baked sand mold to form a combined casting mold. The chills, through their high thermal conductivity, accelerate heat dissipation in specific areas of the casting during the casting process, thus regulating the solidification temperature field distribution.

[0062] The aluminum-silicon alloy melt is cast by controlling the gas pressure inside the pressure vessel, allowing the molten metal to rise smoothly along the riser pipe and fill the mold cavity under pressure. The pressure regime keeps the molten metal in a laminar flow state during the filling process, reducing turbulent air entrainment. During the solidification stage, the continuous back pressure causes the remaining liquid metal between the dendrites to flow towards the shrinkage region, achieving dense feeding and thus obtaining a dense cast blank.

[0063] Finally, surface preparation removes the gating system and surface defects, resulting in castings with acceptable dimensions. Subsequent solution treatment dissolves the strengthening phases in the alloy into the aluminum matrix, forming a supersaturated solid solution. Water quenching, with its rapid cooling rate, preserves the supersaturated state at high temperatures to room temperature, inhibiting the precipitation of equilibrium phases. Aging treatment promotes the decomposition of the supersaturated solid solution, precipitating fine, dispersed strengthening phases. This precipitation strengthening mechanism enhances the mechanical properties of the alloy, ultimately yielding a cast aluminum-silicon alloy with excellent overall performance.

[0064] In some embodiments, surface preparation includes sequential vibration shakeout, sawing, and shot blasting. Vibration shakeout and sawing remove the 3D-printed sand mold and gating system, yielding a rough casting. Shot blasting removes surface sand and oxide scale, while the compressive stress layer introduced by the shot blasting improves the fatigue performance of the casting. In some embodiments, non-destructive testing and repair welding are performed after aging treatment as the final quality assurance step, ensuring the internal quality and surface integrity of the finished casting and ultimately obtaining a qualified finished casting that meets service requirements.

[0065] In a preferred embodiment, in step S2, the sand mold raw material includes silica sand, resin, and curing agent; wherein the average particle size of the silica sand is 90-150 mesh; the amount of resin added is 10-30 mL / Kg; the amount of curing agent added is 4-10 mL / Kg; the resin includes one or more of furan resin, phenolic resin, and urethane resin; the curing agent includes one or more of acid curing agents, gas curing agents, and peroxide curing agents; and / or, the 3D printing thickness is 0.22-0.4 mm, the sand laying speed is 8-15 s / layer, and the gas emission is 8-12 mL / g; and / or, the first baking temperature is 300-400℃, and the time is 1-3 h; and / or, the second baking temperature is 150-200℃, and the time is 1-3 h; and / or, the chilling coating includes one or more of graphite powder coating, zircon powder coating, and chromite powder coating; and / or, the spraying amount of the chilling coating is 1.0-1.8 g / m 2 In 3D printing, the printing thickness refers to the thickness of each printing layer.

[0066] The aforementioned sand particle size range and resin addition amount settings make it easier for the printed sand mold to withstand the filling pressure of differential pressure casting, while also ensuring the mold's permeability to allow gas to escape. The aforementioned gas emission control range further promotes the smooth discharge of gases generated by resin decomposition during differential pressure casting, making it easier to suppress molten metal entrapment and thus reducing the risk of porosity defects. The aforementioned printing thickness and sand laying speed settings further improve the balance between molding accuracy and printing efficiency, making it easier to achieve precise one-time forming of complex internal cavity structures.

[0067] The first baking process on the chills is more effective in removing moisture adsorbed inside the chills and in the sand layer, preventing moisture from vaporizing and forming pores inside the casting during the pouring of high-temperature molten aluminum. Simultaneously, preheating the chills makes the overall temperature distribution of the mold more uniform, further reducing the thermal stress between the chills and the sand mold. The second baking process on the sand mold makes it easier to remove residual volatiles and moisture from the 3D printed sand mold, controlling the gas generation behavior of the sand mold and further promoting the smooth escape of gas during the filling process. Appropriate preheating can also further reduce the cooling rate of the molten metal during filling, improving fluidity and facilitating the complete filling of complex thin-walled structures. Placing cast iron in areas with large wall thickness differences and multiple hot spots allows the chills to absorb local heat from the casting and accelerate cooling, further enhancing the ability to control the solidification sequence and cooling rate in different areas of the casting. This makes it easier for shrinkage porosity to transfer towards the riser, while also refining the local grain structure. For Al-Si alloys, chills can be used to cool at an appropriate cooling rate during casting in areas with large wall thickness differences and multiple hot spots. This can further improve the refinement of α-Al dendrites and eutectic silicon structures, making it easier to obtain uniform precipitation structures in subsequent heat treatment, thereby optimizing the quality of the as-cast substrate.

[0068] In a preferred embodiment, in step S3, the temperature of differential pressure casting is 675~695℃; and / or, the synchronous pressure of differential pressure casting is 0.5~0.65MPa; and / or, step S3 further includes a step of ultrasonically treating the cast aluminum-silicon alloy melt simultaneously with differential pressure casting; the power of ultrasonic treatment is 1~2.5KW; and / or, step S3 includes: placing the cast aluminum-silicon alloy melt in the holding furnace of the differential pressure casting machine, connecting the combined mold and the cast aluminum-silicon alloy melt through a riser pipe, and compressing air through a pressure tank to form pressure on the surface of the cast aluminum-silicon alloy melt, so that the cast aluminum-silicon alloy melt sequentially undergoes rising, filling, pressurizing and holding, and depressurizing along the riser pipe from bottom to top.

[0069] The aforementioned settings for the synchronous pressure range, liquid rise rate range, and filling rate range further enhance the stability of the molten metal rise and filling process, making it easier for the molten aluminum to fill the mold cavity in a laminar flow manner. This minimizes turbulent air entrapment during the filling process and ensures the orderly discharge of gas from the mold cavity. A higher liquid rise rate is more conducive to the smooth rise of the molten aluminum, further reducing impact air entrapment and minimizing excessive cooling of the front-end molten metal due to a slow liquid rise. The aforementioned filling rate range further promotes the laminar flow of molten aluminum into the mold cavity, minimizing turbulent air entrapment during the filling process and facilitating the orderly discharge of gas from the mold cavity.

[0070] Ultrasonic treatment during the filling process can further enhance the cavitation and acoustic flow effects generated by ultrasound in the molten metal, allowing them to play multiple roles. The instantaneous high temperature and pressure generated by the collapse of cavitation bubbles can further break up the already crystallized primary dendrites, making it easier to achieve grain refinement and second-phase fragmentation. The forced convection at both the macroscopic and microscopic scales generated by the acoustic flow effect can further improve the uniformity of the temperature and concentration fields, making it easier to suppress solute segregation. Ultrasound can further promote the wetting and adhesion of the molten metal to the sand mold wall, thereby improving the surface quality of the casting. Applying ultrasound simultaneously during differential pressure filling makes it easier to achieve a more direct microstructure control effect than pre-ultrasonic treatment in the ladle, allowing the ultrasonic effect to run through the entire filling and initial solidification process, thus optimizing the casting microstructure throughout the entire process.

[0071] In a preferred embodiment, in step S3, the rate of liquid influx is 1.1~1.5 kPa / s, the rate of filling is 0.8~1.1 kPa / s; and / or, the pressurization and holding stage includes a pressurization and shelling stage and a holding and crystallization stage. The pressurization and shelling stage has a pressurization rate of 1.0~1.2 kPa / s, a pressurization value of 12~18 kPa, and a holding time of 8~15 s. The holding and crystallization stage has a pressurization value of 0.4~0.5 MPa and a holding time of 300~600 s.

[0072] The aforementioned ranges for shell formation pressure and time further promote the formation of a thin, dense solidified shell layer between the molten metal in the mold cavity and the sand mold wall under pressure. This shell layer better protects the surface quality of the casting and provides structural support for subsequent high-pressure feeding. The aforementioned ranges for crystallization pressure and holding time further enhance the pressure effect on the feeding channels during solidification, making it easier to force the interdendritic liquid phase to flow towards the shrinkage porosity region, thereby further eliminating shrinkage porosity and shrinkage defects and resulting in a completely dense microstructure in the casting. Differential pressure casting within the aforementioned parameter range further improves the controllability of the filling speed and the flexibility of the gate position, making it easier for gas to be discharged from the mold cavity in an orderly manner. Simultaneously, the forced feeding effect during the crystallization pressure stage further increases the density of the casting, thus making it easier for the alloy's mechanical properties to be fully realized.

[0073] In a preferred embodiment, step S4, the solution treatment includes a first solution treatment, a second solution treatment, and a third solution treatment performed sequentially; the temperature of the first solution treatment is 360~380℃, and the holding time is 8~14h; the temperature of the second solution treatment is 490~510℃, and the holding time is 3~5h; the temperature of the third solution treatment is 515~535℃, and the holding time is 8~12h; and / or, the water quenching temperature is 60~80℃, and the quenching transfer time is 5~10s; and / or, the aging treatment temperature is 160~175℃, and the holding time is 10~14h.

[0074] The aforementioned first solution temperature range facilitates the complete dissolution and uniform distribution of Sn, while simultaneously promoting the initial diffusion and homogenization of elements such as Mg and Si, thus preparing the composition and microstructure for subsequent high-temperature processing. This temperature range further enhances the ability to suppress the local enrichment of low-melting-point elements such as Mg and Sn, thereby making it easier to prevent overheating during subsequent high-temperature stages.

[0075] The aforementioned setting of the second solid solution temperature range and holding time range can further enhance the dissolution effect of major strengthening elements such as Cu and Mg into the aluminum matrix, promote the gradual re-dissolution of the precursors of the θ and Q phases, and make it easier for eutectic silicon to begin granulation transformation. The setting of the holding time range in this stage can further inhibit the premature excessive precipitation of elements such as Zr, V, and Ti in this temperature range, thereby making it easier to retain sufficient supersaturation for the uniform nucleation of nanophases in the subsequent high-temperature stage.

[0076] The aforementioned third solution temperature range and holding time range can further enhance the granulation degree of eutectic silicon, making it easier to push the solid solubility of elements such as Cu and Mg to their maximum value. Since the first two stages of solution treatment have basically eliminated the local segregation of the low-melting-point phase, the above temperature range settings can further enhance the control of overheating risk, thus making it easier to effectively suppress the risk of overheating of the low-melting-point phase and premature coarsening of the high-temperature phase while ensuring that the eutectic silicon is fully spheroidized and the solute elements are fully dissolved.

[0077] The aforementioned water quenching temperature range and transfer time range further enhance the cooling rate of supersaturated solid solutions after high-temperature solution treatment, making it easier to rapidly freeze the microstructure and thus better preserve the maximum supersaturation for subsequent aging precipitation. The aforementioned water quenching temperature range also helps to mitigate thermal stress during the quenching process, making it easier to suppress deformation or microcracks caused by excessive thermal stress, while ensuring a sufficient cooling rate. The aforementioned rapid transfer time range further enhances the ability to suppress premature precipitation of solute atoms during the transfer process.

[0078] The aforementioned aging treatment temperature and holding time ranges can further enhance the precipitation driving force and growth kinetics of the θ′ phase, Q phase, and Al3(Zr,V,Ti) nanophase. The θ′ phase is more likely to precipitate in a coherent, disk-like manner, thus providing the greatest contribution to precipitation strengthening; the Q phase is more likely to precipitate uniformly in a fine-particle form, which is more conducive to maintaining high-temperature strength; the Al3(Zr,V,Ti) phase is more likely to further precipitate and grow during aging, forming a high-temperature stable network. The synergistic effect of these three phases can further improve the matching effect between the alloy's strength and plasticity.

[0079] In a preferred embodiment, step S1 includes: step S11, melting and holding aluminum ingots at a first holding temperature, sequentially adding AlSi master alloy, AlCu master alloy, AlNi master alloy, AlMn master alloy, AlV master alloy and Sn ingot at the first holding temperature to obtain a first intermediate melt; step S12, holding the first intermediate melt at a second holding temperature, sequentially adding AlZr master alloy and Mg ingot at the second holding temperature to melt, then adding AlTi master alloy and mixing to obtain a second intermediate melt; passing argon gas through the second intermediate melt for refining, adding a refining agent to obtain a third intermediate melt; adding AlSr master alloy to the third intermediate melt for modification treatment to obtain a cast aluminum-silicon alloy melt.

[0080] This invention involves the stepwise addition of raw materials to the melt. First, high-melting-point master alloys such as AlSi, AlCu, AlNi, AlMn, and AlV are added for smelting. This facilitates the formation of stable intermetallic compounds and ensures the complete dissolution and uniform distribution of these high-melting-point elements. Sn ingots, with their lower melting point, are added together with the aforementioned high-melting-point master alloys. This allows Sn to fully dissolve and uniformly distribute in the high-temperature melt, further promoting the "vacancy engineering" effect of Sn and its uniform distribution as a heterogeneous nucleation core.

[0081] The subsequent addition of AlZr and Mg is primarily due to the limited solubility and high melting point of Zr in molten aluminum. In this invention, AlZr is added at the second holding temperature to reduce Zr supersaturation and the driving force for primary phase precipitation, allowing Zr to enter the matrix in a solid solution state. This is more conducive to the uniform precipitation of nano-sized Al3(Zr,V,Ti) phases during subsequent aging, thus more easily suppressing the formation of insoluble primary Al3Zr coarse particles and preventing these coarse particles from adversely affecting the alloy's plasticity and fatigue properties. Simultaneously, because Mg is chemically reactive and has a low melting point, shortening the residence time of Mg in the melt effectively reduces Mg oxidation loss in the high-temperature melt and volatilization loss due to high vapor pressure, making the alloy composition more accurately meet design values ​​and further reducing the hydrogen content of the melt. Finally, AlTi is added. Based on the promoting effect of Ti on the nucleation of aging precipitation phases, it further promotes the dispersion of Ti in the form of fine Al3Ti particles during solidification and subsequent heat treatment, so as to provide the dual role of grain refinement and Al3(Zr,V,Ti) nanophase nucleation, further reducing the coarsening or precipitation of Ti due to long-term high-temperature holding, thereby maintaining its effective nucleation promoting effect.

[0082] In addition, the addition of Sr for modification treatment can change the growth mode and twin density of eutectic silicon, transforming coarse lamellar or needle-like eutectic silicon into fine fibrous or coral-like structures, which can further improve the morphological distribution of eutectic silicon. After Sr modification, the cutting effect of eutectic silicon on the aluminum matrix is ​​greatly reduced, making it easier for the alloy to maintain good elongation while obtaining higher strength.

[0083] In a preferred embodiment, in step S11, the weight percentage of Si in the AlSi master alloy is 20-30%; and / or, the weight percentage of Cu in the AlCu master alloy is 30-50%; and / or, the weight percentage of Mn in the AlMn master alloy is 10-20%; and / or, the weight percentage of Zr in the AlZr master alloy is 5-10%; and / or, the weight percentage of V in the AlV master alloy is 5-10%; and / or, the weight percentage of Ti in the AlTi master alloy is 10-15%; and / or, the weight percentage of Ni in the AlNi master alloy is 10-20%; and / or, the melting temperature is 730-750°C, and the melting time is 0.5-1 hour. In some embodiments, the crucible used for melting is made of cast iron, and its inner wall is pre-coated with a water-based cast iron crucible coating; the coating operation is performed when the crucible temperature reaches 50-80°C.

[0084] In a preferred embodiment, the first holding temperature is 720~740℃; and / or, the second holding temperature is 700~720℃; and / or, step S11 further includes the following steps: before melting, aluminum ingots, AlSi master alloys, AlCu master alloys, AlMn master alloys, AlZr master alloys, AlV master alloys, AlTi master alloys, AlNi master alloys, Sn ingots, and Mg ingots are respectively subjected to a third baking; the third baking temperature is 150~250℃, and the time is 1~3h; and / or, in step S12, the refining temperature is 730~740℃, the time is 20~30min, and the rotation speed is 400~600r / min; and / or, the refining agent includes one or more of AlTiB, AlTiC, and TCB, and the weight ratio of the refining agent to the first intermediate melt is 0.3~0.8%; and / or, the modification treatment temperature is 720~740℃, and the time is 20min~40min.

[0085] Hydrogen in molten aluminum alloys is the main source of pinhole defects, and baking to remove moisture helps to further reduce the pinhole rate. The aforementioned third baking treatment can remove moisture and oil adsorbed on the surface of the raw materials, prevent the introduction of hydrogen atoms by water vapor decomposition during the smelting process, and control the gas content of the alloy from the source.

[0086] The aforementioned refining conditions facilitate the adsorption and carry of hydrogen atoms from the melt to the surface by argon bubbles during their ascent, utilizing the principle of partial pressure difference, thereby further reducing the hydrogen content of the melt. The stirring effect of the rotating gas flow further promotes the collision and aggregation of non-metallic inclusions, causing them to be carried away with the rising bubbles, thus achieving melt purification. Simultaneously, it further improves degassing and impurity removal efficiency, making it easier to suppress secondary oxidation and gas entrainment.

[0087] The addition of the aforementioned grain refiner can further refine the α-Al grain size. The refined equiaxed grain structure facilitates a shorter microsegregation distance of solute elements, resulting in a more uniform distribution of transition elements such as Zr, V, and Ti. This is more conducive to the uniform precipitation of subsequent nanophases and also makes it easier to reduce the degree and size of the second phase at grain boundaries, thereby improving alloy plasticity. It can further improve the feeding efficiency of castings, thus reducing shrinkage defects. The synergistic effect of grain refinement and differential pressure casting can further improve the density of castings, making it easier to obtain a fine-grained, dense, and high-quality microstructure. The elements introduced by the grain refiner are considered in the proportioning, but this is not a simple linear addition. Instead, a post-furnace addition method is used: the matrix elements are proportioned to near the lower limit of the design before the furnace, reserving a buffer for fluctuations in the yield and ensuring that the final composition does not reach the upper limit. This provides sufficient margin for the process window, ensuring that the alloy composition always remains within the design range, thus preventing the alloy composition from exceeding the design range of this application.

[0088] Typical, but not limiting, weight ratios of Cu to Mg are 4:1, 4.2:1, 4.4:1, 4.6:1, 4.8:1, 5:1, 5.2:1, 5.4:1, 5.6:1, 5.8:1, 6:1, or any range of two of these values; weight ratios of Zr, V, and Ti are 0.4:0.6:1, 0.5:0.7:1, 0.6:0.8:1, 0.7:0.9:1, 0.8:1.0:1, 1:1.1:1, 1.2:1.2:1, 1.5:1.3:1, 1.6:1.4:1, 1.8:1.5:1, or any range of two of these values.

[0089] Typical, but not limiting, weight ratios of Ni to Sn are 4:1, 4.6:1, 5.2:1, 5.8:1, 6.4:1, 7:1, 7.6:1, 8.2:1, 8.8:1, 9.4:1, 10:1, or any range of any two of these values; the ratio of the sum of the weights of Zr, V, and Ti to the sum of the weights of Ni and Sn is 1.72:1, 2.14:1, 2.56:1, 2.98:1, 3.4:1, 3.82:1, 4.24:1, 4.66:1, 5:1, or any range of any two of these values.

[0090] Typically, but not limitingly, in step S11, the temperature of the third baking is 150°C, 161°C, 172°C, 183°C, 194°C, 205°C, 216°C, 227°C, 238°C, 250°C, or a range of any two of these values; the baking time is 1h, 1.2h, 1.4h, 1.6h, 1.8h, 2.0h, 2.2h, 2.4h, 2.6h, 3h, or a range of any two of these values.

[0091] Typical, but not limiting, temperatures for the deterioration treatment are 720℃, 722℃, 724℃, 726℃, 728℃, 730℃, 732℃, 734℃, 736℃, 740℃, or any two of these values; and deterioration treatment times are 20 min, 22 min, 24 min, 26 min, 28 min, 30 min, 32 min, 34 min, 36 min, 40 min, or any two of these values.

[0092] Typical, but not limiting, 3D printing thicknesses are 0.22mm, 0.24mm, 0.26mm, 0.28mm, 0.30mm, 0.32mm, 0.34mm, 0.36mm, 0.38mm, 0.4mm, or any two of these values; sand-laying speeds are 8s / layer, 9s / layer, 10s / layer, 11s / layer, 12s / layer, 13s / layer, 14s / layer, 15s / layer, or any two of these values; and gas generation is 8mL / g, 9mL / g, 10mL / g, 11mL / g, 12mL / g, or any two of these values.

[0093] The present application will be further described in detail below with reference to specific embodiments, which should not be construed as limiting the scope of protection claimed in the present application.

[0094] Example 1

[0095] Step S11: Prepare the raw materials according to the composition ratio of the cast aluminum-silicon alloy (see Tables 1 to 2). The smelting raw materials specifically include high-purity aluminum, AlSi master alloy, AlCu master alloy, AlMn master alloy, AlZr master alloy, AlV master alloy, AlTi master alloy, AlNi master alloy, pure Sn, and pure Mg. After the above raw materials are baked for the third time in the raw material oven, the temperature of the third baking is 200℃ and the time is 2h. The composition parameters of each master alloy are shown in Table 3.

[0096] High-purity aluminum is placed in a crucible for sequential melting and first holding. AlSi master alloy, AlCu master alloy, AlNi master alloy, AlMn master alloy, AlV master alloy and pure Sn are added sequentially at the first holding temperature to obtain the first intermediate melt.

[0097] Step S12: The first intermediate melt is subjected to a second heat treatment. At the second heat treatment temperature, AlZr master alloy and pure Mg are added sequentially and melted. Then, AlTi master alloy is added and mixed to obtain the second intermediate melt. Argon gas is introduced into the second intermediate melt for refining to obtain the third intermediate melt. The second intermediate melt is refined by a refining degasser with argon gas. After refining, slag is removed. After the composition is qualified, 0.4% of a refining agent (AlTiB) is added to obtain the third intermediate melt. Specific parameters are shown in Tables 4 and 5. AlSr master alloy is added to the third intermediate melt and modified at 730℃ for 30 minutes to obtain a cast aluminum-silicon alloy melt. Specific parameters are shown in Table 4.

[0098] Step S2: Select the target casting, use casting simulation software to design and optimize the design of the gating system, risers and chills, simulate and predict casting defects, optimize the casting process, and realize the design of the sand mold 3D printing model. Then, use sand mold 3D printing equipment to 3D print the sand mold material to obtain the sand mold. Among them, the resin used in the sand mold material is phenolic resin, and the curing agent is an acid curing agent. Other specific parameters are shown in Table 4.

[0099] A casting mold is set up for areas with large wall thickness differences and multiple hot spots, and cast iron is used to cast the chills.

[0100] The chills were sequentially subjected to sand coating and a first baking process to obtain chills after the first baking. A chilling coating (zircon powder coating, spraying amount 1.2 g / m²) was then sprayed onto the chills and the sand mold. 2 After that, a second baking is performed to obtain a chill after the second baking and a sand mold after baking; the chill after the second baking and the sand mold after baking are assembled to obtain a combined mold; specific parameters are shown in Table 5;

[0101] Step S3: Place the molten aluminum-silicon alloy in the holding furnace of the differential pressure casting machine. Connect the combined mold and the molten aluminum-silicon alloy through a riser pipe. Compressed air from a pressure tank creates pressure on the surface of the molten aluminum-silicon alloy, causing the molten aluminum-silicon alloy to rise, fill, pressurize, and depressurize sequentially from bottom to top along the riser pipe. Simultaneously, perform polar ultrasonic treatment on the molten aluminum-silicon alloy to obtain the casting blank. Specific parameters are shown in Table 6.

[0102] Step S4: Vibrate the casting blank to remove sand, saw off the riser and gating system, then shot blast to remove surface burrs, and machine to the specified dimensions to obtain the casting; the casting is then subjected to solution treatment, water quenching and aging treatment in sequence to obtain the cast aluminum-silicon alloy; specific parameters are shown in Table 7.

[0103] Examples 2 to 8

[0104] The difference from Example 1 is that the composition ratio of the cast aluminum-silicon alloy and the casting raw materials are different, as shown in Tables 1 to 3; the parameters of each step are different, as shown in Tables 4 to 7.

[0105] Example 9

[0106] The difference from Example 1 is that,

[0107] In step S11, the melting temperature is 730℃ and the time is 1 hour;

[0108] In step S12, the refining agent used is AlTiC, and the amount added is 0.3% of the first intermediate melt; the temperature of the modification treatment is 720℃ and the time is 40min.

[0109] In step S2, the resin used is furan resin, and the curing agent is a gas curing agent; the chilling coating used is graphite powder coating, and the spraying amount is 1.0 g / m². 2 .

[0110] Example 10

[0111] The difference from Example 1 is that,

[0112] In step S11, the melting temperature is 750℃ and the melting time is 0.5h;

[0113] In step S12, the refining agent used is TCB, and the amount added is 0.8% of the first intermediate melt; the temperature of the modification treatment is 740℃ and the time is 20min.

[0114] In step S2, the resin used is urethane resin, and the curing agent is a peroxide curing agent; the chilling coating used is chromite powder coating, and the spraying amount is 1.0 g / m². 2 .

[0115] Comparative Examples 1 to 3

[0116] The difference from Example 1 is that the composition ratio of the cast aluminum-silicon alloy and the casting raw materials are different, as shown in Tables 1 to 3; the parameters of each step are different, as shown in Tables 4 to 7.

[0117] Performance testing:

[0118] The cast aluminum-silicon alloys obtained in the above embodiments and comparative examples were analyzed and tested as follows, and the results are shown in Table 8.

[0119] 1. Tensile Strength: The room temperature tensile test shall be conducted according to GB / T 228.1-2021 Metallic materials, tensile testing—Part 1: Test method at room temperature, on an electronic universal testing machine, using a standard dumbbell-shaped specimen. The high temperature tensile test shall be conducted according to GBT-228.2-2015 Metallic materials, tensile testing—Part 2: Test method at high temperature, on an electronic universal testing machine, using a standard dumbbell-shaped specimen.

[0120] 2. Microstructure Analysis: The finished casting was analyzed using an EBSD scanning electron microscope, magnified 200x along the casting direction. The average grain size was measured using OIM8.0 analysis software to evaluate the grain size of the casting samples.

[0121] 3. Non-destructive testing: Non-destructive testing of castings shall be carried out in accordance with GB / T 11346-2018 Radiographic testing of aluminum alloy castings and evaluated in accordance with HB 963-2021 Specification for aluminum alloy castings.

[0122] Table 1

[0123]

[0124] Table 2

[0125]

[0126] Table 3

[0127]

[0128] Table 4

[0129]

[0130] Table 5

[0131]

[0132] Table 6

[0133]

[0134] Table 7

[0135]

[0136] Table 8

[0137]

[0138] The EBSD microstructure of the alloy in Example 1 under the T6 condition is shown in the figure. Figure 1 As can be seen, the alloy matrix consists of fine equiaxed crystals with small and uniform grain size, indicating that the alloy achieved a good recrystallized structure after multi-stage solution treatment and aging. The metallographic (OM) image of the alloy in the as-cast state in Example 1 is shown below. Figure 2 As can be seen, the internal structure of the alloy is uniformly distributed, and the eutectic silicon phase, after modification treatment, appears as fine fibers or granules without obvious coarse defects, demonstrating the effect of differential pressure casting on improving the density of the casting structure. A transmission electron microscope (TEM) image of the alloy in Example 1 under the T6 condition is shown below. Figure 3 Needle-shaped precipitates, Q phase, are uniformly dispersed in the matrix. The transmission electron microscopy diffraction pattern of the alloy in Example 1 under the T6 condition is shown below. Figure 4 Clear diffraction spots and superlattice spots are visible, further confirming that the precipitated phase is a Q phase with a specific crystal structure, which can verify the phenomenon that the multi-element microalloying design induces the formation of the above-mentioned high thermal stability strengthening phase during the aging process.

[0139] As can be seen, in Comparative Example 1, insufficient raw material baking temperature led to an increase in the gas content of the melt. Furthermore, insufficient addition of sand mold resin and curing agent, as well as insufficient chill baking temperature and sand mold baking temperature, resulted in severe gas absorption in the overall casting. Additionally, insufficient filling pressure and low filling rate during mold filling also led to insufficient filling and feeding, resulting in casting quality failing to meet overall requirements. In Comparative Example 2, the smaller mesh size of the 3D printing sand resulted in a lower precision grade of the cast part. Furthermore, excessively high chill baking temperature and sand mold baking temperature led to poor sand adhesion on the chill and damaged the sand mold structure, causing a series of problems such as decreased strength, cracking, and deformation, ultimately leading to casting defects such as sand adhesion and porosity. Excessively fast liquid rise and filling speeds caused turbulence and air entrapment, obstructed cavity venting, and erosion of the cavity and coating, ultimately reducing the density of the casting. In Comparative Example 3, the excessive addition of resin and curing agent significantly increases gas generation, leading to porosity in the casting. It also reduces the strength and dimensional accuracy of the sand mold, deteriorates the filling capacity, and causes defects such as slag inclusions. The process requires precise control of their proportions. In addition, the casting uses single-stage solid solution, resulting in insufficient solid solution of some elements, which makes it impossible to obtain castings that meet performance requirements during subsequent aging treatment.

[0140] As can be seen from the above, this invention, through the use of multi-element microalloying components such as Si, Cu, Mn, Mg, Zr, V, Ti, Ni, Sn, Fe, and Sr, achieves significant refinement of grain size. The resulting cast aluminum-silicon alloy exhibits comprehensive improvements in room temperature tensile strength, elongation, and high temperature tensile strength and elongation, giving it high room temperature strength, high toughness, and excellent high temperature strength. Compared with traditional cast Al-Si alloys, this invention, through comprehensive innovation in multi-element microalloying composition design, differential pressure casting and ultrasonic treatment synergistic process, and multi-stage heat treatment regime, achieves significant improvement in the mechanical properties of Al-Si alloys. It has significant industrial practical value and broad application prospects, and is especially suitable for high-end fields such as aerospace, weaponry, and high-speed rail.

[0141] Furthermore, it can be seen that the overall effect is better when all process parameters are within the preferred range of this invention. Specifically, in terms of composition design, Al is formed by synergistically adding Zr, V, and Ti in a defined ratio to achieve the desired effect. 21.4 Si 4.1 Ti 3.5 VZr 3.9The addition of multiple high-thermal-stability, multi-component heat-resistant phases is more conducive to effectively pinning grain boundaries and hindering high-temperature dislocation movement. The synergistic addition of Cu and Mg in the aforementioned weight ratio, along with Ni and Sn in the same weight ratio, allows precipitation-strengthening phases such as θ'(Al2Cu) and Q(AlCuMgSi) to work together with the high-temperature stable Ni-containing phase, further balancing precipitation strengthening and high-temperature stability. In terms of manufacturing process, differential pressure casting combined with ultrasonic treatment within the aforementioned range during the filling process is employed. Nonlinear pressurization further controls the filling process, and the ultrasonic cavitation effect further refines the eutectic silicon and iron-rich phases, thus reducing casting defects such as shrinkage porosity and improving density. Combined with sand mold 3D printing precision forming technology, high-precision manufacturing of complex castings can be achieved. For multi-stage heat treatment, a three-stage solution treatment within the aforementioned temperature range, combined with aging treatment within the aforementioned range, further reduces overheating and fully dissolves the strengthening phases. The aging process further enhances the effect of dispersion precipitation strengthening, ultimately giving the casting body both excellent room-temperature mechanical properties and high-temperature heat resistance.

[0142] 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 cast aluminum-silicon alloy, characterized in that, By weight percentage, the cast aluminum-silicon alloy comprises 7.6-10% Si, 1.0-2.0% Cu, 0.1-0.3% Mn, 0.2-0.4% Mg, 0.052-0.18% Zr, 0.075-0.165% V, 0.1-0.15% Ti, 0.05-0.2% Ni, 0.01-0.05% Sn, 0.01-0.05% Fe, 0.01-0.05% Sr, with the balance being Al and unavoidable impurity elements; wherein the total amount of the impurity elements is ≤0.12%.

2. The cast aluminum-silicon alloy according to claim 1, characterized in that, The weight ratio of Cu to Mg is (4~6):1; and / or, The weight ratio of Zr, V, and Ti is (0.4~1.8):(0.6~1.5):1; and / or, The weight ratio of Ni to Sn is (4~10):1; and / or, The ratio of the weight sum of Zr, V and Ti to the weight sum of Ni and Sn is (1.72~5):

1.

3. The cast aluminum-silicon alloy according to claim 1 or 2, characterized in that, By weight percentage, the cast aluminum-silicon alloy comprises 8-9.5% Si, 1.2-1.8% Cu, 0.13-0.25% Mn, 0.25-0.35% Mg, 0.08-0.16% Zr, 0.078-0.15% V, 0.1-0.15% Ti, 0.08-0.18% Ni, 0.02-0.05% Sn, 0.01-0.05% Fe, 0.01-0.04% Sr, with the balance being Al and unavoidable impurity elements.

4. The method for preparing the cast aluminum-silicon alloy according to any one of claims 1 to 3, characterized in that, Includes the following steps: Step S1: According to the composition ratio of the cast aluminum-silicon alloy, the raw materials are batched, smelted and refined in sequence to obtain the cast aluminum-silicon alloy melt; Step S2: 3D print the sand mold material to obtain a sand mold; set up a casting model for areas with large wall thickness differences and multiple hot spots, and cast it using cast iron or cast aluminum to obtain a chill; sequentially coat the chill with sand and bake it for the first time to obtain a chill after baking; spray the chill after baking and the sand mold with a chilling coating, and then bake it for the second time to obtain a chill after baking and a baked sand mold. The second baked chill is assembled with the baked sand mold to obtain a combined casting mold; Step S3: The cast aluminum-silicon alloy melt is subjected to differential pressure casting in the combined mold to obtain a cast blank. Step S4: Surface straightening of the casting blank to obtain the casting; The casting is subjected to solution treatment, water quenching and aging treatment in sequence to obtain the cast aluminum-silicon alloy.

5. The method for preparing cast aluminum-silicon alloy according to claim 4, characterized in that, In step S2 The sand molding raw materials include silica sand, resin, and curing agent; wherein... The average particle size of the silica sand is 90-150 mesh; the amount of resin added is 10-30 mL / kg; the amount of curing agent added is 4-10 mL / kg; the resin includes one or more of furan resin, phenolic resin, and urethane resin; the curing agent includes one or more of acid curing agents, gas curing agents, and peroxide curing agents; and / or, The 3D printing process has a printing thickness of 0.22~0.4mm, a sand-laying speed of 8~15s / layer, and a gas emission rate of 8~12 mL / g; and / or, The first baking temperature is 300~400℃, and the time is 1~3 hours; and / or, The second baking temperature is 150~200℃, and the time is 1~3 hours; and / or, The chilling coating includes one or more of graphite powder coating, zircon powder coating, and chromite powder coating; and / or, The spraying amount of the chilling coating is 1.0~1.8 g / m³. 2 .

6. The method for preparing cast aluminum-silicon alloy according to claim 4 or 5, characterized in that, In step S3 The differential pressure casting temperature is 675~695℃; and / or, The synchronous pressure of the differential pressure casting is 0.5~0.65MPa; and / or, Step S3 further includes the step of ultrasonically treating the cast aluminum-silicon alloy melt simultaneously with the differential pressure casting; the power of the ultrasonic treatment is 1~2.5KW; And / or, Step S3 includes: placing the cast aluminum-silicon alloy melt in the holding furnace of the differential pressure casting machine, connecting the combined mold and the cast aluminum-silicon alloy melt through a riser pipe, and compressing air through a pressure tank to create pressure on the surface of the cast aluminum-silicon alloy melt, so that the cast aluminum-silicon alloy melt sequentially rises, fills the mold, pressurizes, holds pressure, and depressurizes along the riser pipe from bottom to top.

7. The method for preparing cast aluminum-silicon alloy according to claim 6, characterized in that, In step S3 The rate of liquid inflow is 1.1~1.5 kPa / s, and the rate of mold filling is 0.8~1.1 kPa / s; and / or, The pressurization and holding stage includes a pressurization and crust formation stage and a holding and crystallization stage. The pressurization rate of the pressurization and crust formation stage is 1.0~1.2 kPa / s, the pressurization value is 12~18 kPa, and the time is 8~15 s. The pressurization value of the holding and crystallization stage is 0.4~0.5 MPa, and the holding time is 300~600 s.

8. The method for preparing cast aluminum-silicon alloy according to claim 4, characterized in that, In step S4 The solution treatment includes a first solution treatment, a second solution treatment, and a third solution treatment performed sequentially; the temperature of the first solution treatment is 360~380℃, and the holding time is 8~14h; the temperature of the second solution treatment is 490~510℃, and the holding time is 3~5h; the temperature of the third solution treatment is 515~535℃, and the holding time is 8~12h; and / or, The water quenching temperature is 60~80℃, and the quenching transfer time is 5~10s; and / or, The aging treatment temperature is 160~175℃, and the holding time is 10~14h.

9. The method for preparing cast aluminum-silicon alloy according to claim 4, characterized in that, Step S1 includes: Step S11: The aluminum ingot is smelted and subjected to a first heat preservation. At the first heat preservation temperature, AlSi master alloy, AlCu master alloy, AlNi master alloy, AlMn master alloy, AlV master alloy and Sn ingot are added in sequence to obtain a first intermediate melt. Step S12: The first intermediate melt is subjected to a second heat preservation. At the second heat preservation temperature, AlZr master alloy and Mg ingot are added sequentially and melted. Then, AlTi master alloy is added and mixed to obtain a second intermediate melt. Argon gas is introduced into the second intermediate melt for refining. A refining agent is added to obtain a third intermediate melt. AlSr master alloy is added to the third intermediate melt for modification treatment to obtain the cast aluminum-silicon alloy melt.

10. The method for preparing cast aluminum-silicon alloy according to claim 9, characterized in that, In step S11 The AlSi master alloy contains 20-30% Si by weight; and / or, the AlCu master alloy contains 30-50% Cu by weight; and / or, the AlMn master alloy contains 10-20% Mn by weight; and / or, the AlZr master alloy contains 5-10% Zr by weight; and / or, the AlV master alloy contains 5-10% V by weight; and / or, the AlTi master alloy contains 10-15% Ti by weight; and / or, the AlNi master alloy contains 10-20% Ni by weight; and / or, The melting temperature is 730~750℃, and the time is 0.5~1h; and / or, The temperature of the first insulation is 720~740℃; and / or, The second insulation temperature is 700~720℃; and / or, Step S11 further includes the following steps: before the melting, the aluminum ingot, the AlSi master alloy, the AlCu master alloy, the AlMn master alloy, the AlZr master alloy, the AlV master alloy, the AlTi master alloy, the AlNi master alloy, the Sn ingot, and the Mg ingot are each subjected to a third baking; the temperature of the third baking is 150~250℃, and the time is 1~3h; and / or, In step S12 The refining temperature is 730~740℃, the time is 20~30 min, and the rotation speed is 400~600 r / min; and / or, The refining agent includes one or more of AlTiB, AlTiC, and TCB, and the weight ratio of the refining agent to the first intermediate melt is 0.3~0.8%; and / or, The temperature for the degradation treatment is 720~740℃, and the time is 20min~40min.

Images