High-strength heat-resistant cast aluminum alloy and method for manufacturing the same

Abstract

The application provides a high-toughness heat-resistant cast aluminum alloy and a preparation method thereof. The alloy comprises Cu 4.5-5.5%, Mn 0.4-0.8%, Mg 0.15-0.3%, Zr 0.15-0.25%, V 0.08-0.15%, Ti 0.08-0.15%, B 0.005-0.02%, Cd 0.15-0.25%, Ag 0.05-0.20%, RE 0.05-0.20%, Fe 0.01-0.05%, Si 0.01-0.03%, and the balance of Al and inevitable impurity elements. The application adopts multi-element micro-alloying composition, the tensile strength of the obtained high-toughness heat-resistant cast aluminum alloy at room temperature and high temperature is significantly improved, the elongation is moderate, the yield strength is synchronously improved, the grain size is significantly refined, and the matching of high strength, high toughness and heat resistance is realized.

Description

Technical Field

[0001] This invention relates to the field of alloy materials technology, and more specifically, to a high-strength, high-toughness, heat-resistant cast aluminum alloy and its preparation method. Background Technology

[0002] Cast aluminum alloys are widely used in aerospace, rail transportation, and automotive manufacturing—fields with high requirements for lightweight materials and comprehensive mechanical properties—due to their excellent casting performance, low density, and good corrosion resistance. As high-end equipment develops towards high temperature, high load, and long service life, higher demands are placed on the room temperature strength, plasticity, and high-temperature tensile strength and creep resistance of cast aluminum alloys. Especially in operating environments around 200℃, traditional cast aluminum alloys often exhibit significant strength reduction and plasticity deterioration, making it difficult to meet practical engineering needs.

[0003] Currently, commonly used heat-resistant cast aluminum alloys mainly include Al-Si, Al-Cu, and Al-Mg systems. While Al-Si alloys possess good casting fluidity and corrosion resistance, their high-temperature strength is insufficient, making them unsuitable for load-bearing structural components. Al-Mg alloys exhibit excellent corrosion resistance, but their casting performance is poor, and improvements in high-temperature mechanical properties are limited. Al-Cu alloys (such as ZL201 and ZL205) have high room-temperature strength, but their high-temperature performance remains unsatisfactory, exhibiting problems such as a high tendency for hot cracking and low plasticity. To improve heat resistance, researchers have attempted to add transition elements such as Mn, Zr, V, and Ti, or rare earth elements (RE), to form high-temperature stable dispersed or precipitated phases, suppressing grain boundary slip and dislocation movement.

[0004] However, existing alloy systems still have the following shortcomings. First, it is difficult to balance high-temperature strength and room-temperature plasticity: most existing alloys sacrifice room-temperature plasticity while improving high-temperature strength, especially with generally low elongation (usually below 5%), limiting their application under impact or complex stress conditions. Second, the alloying elements are not properly matched, making it difficult to control precipitates: although some alloys have added elements such as Zr, Ti, and RE, the lack of systematic optimization of element ratios and preparation processes makes it difficult to form fine, uniform, and thermally stable precipitates, resulting in insignificant strengthening effects, or even the appearance of coarse phases or segregation. Third, the control of impurity elements is not strict: common impurities such as Fe and Si easily form brittle needle-like phases in cast aluminum alloys, significantly reducing the alloy's toughness and fatigue performance, and existing technologies lack clear definitions of the content range and synergistic control effects of such impurities.

[0005] Therefore, developing a cast aluminum alloy that combines high-temperature / room-temperature strength, good plasticity (high elongation), and excellent high-temperature (around 200℃) tensile strength has significant engineering value and market prospects. Summary of the Invention

[0006] The main objective of this invention is to provide a high-strength, high-toughness, and heat-resistant cast aluminum alloy and its preparation method, so as to solve the problem that aluminum alloys in the prior art are difficult to achieve simultaneously with high strength, high toughness, excellent high-temperature performance, and good casting performance.

[0007] To achieve the above objectives, according to one aspect of the present invention, a high-strength, high-toughness, and heat-resistant cast aluminum alloy is provided, comprising, by weight percentage: Cu 4.5~5.5%, Mn 0.4~0.8%, Mg 0.15~0.3%, Zr 0.15~0.25%, V 0.08~0.15%, Ti 0.08~0.15%, B 0.005~0.02%, Cd 0.15~0.25%, Ag 0.05~0.20%, RE 0.05~0.20%, Fe 0.01~0.05%, Si 0.01~0.03%, with the balance being Al and unavoidable impurity elements, wherein the individual impurity element is ≤0.03% and the total amount of impurity elements is ≤0.12%.

[0008] Further, RE includes La and Ce, and the weight ratio of La to Ce is 1:(1~3); and / or the weight ratio of Mn to Cu is 0.08~0.15; and / or the weight ratio of Zr to Ti is 1.2~2.0; and / or the weight ratio of Ti to B is 6~26.

[0009] Furthermore, by weight percentage, the high-strength, high-toughness, and heat-resistant cast aluminum alloy comprises Cu 4.8~5.2%, Mn 0.5~0.7%, Mg 0.2~0.28%, Zr 0.18~0.24%, V 0.10~0.14%, Ti 0.09~0.14%, B 0.006~0.015%, Cd 0.18~0.24%, Ag 0.08~0.15%, RE 0.08~0.18%, Fe 0.01~0.05%, Si 0.01~0.03%, with the balance being Al and unavoidable impurity elements.

[0010] According to another aspect of the present invention, a method for preparing the above-mentioned high-strength, high-toughness, and heat-resistant cast aluminum alloy is provided, comprising the following steps: Step S1, batching, melting, and refining the high-strength, high-toughness, and heat-resistant cast aluminum alloy in sequence to obtain an aluminum alloy melt; Step S2, transferring the aluminum alloy melt to a casting machine for casting, the casting including low-pressure casting or differential pressure casting, to obtain a cast blank; Step S3, subjecting the cast blank to two-stage solution treatment, quenching, and two-stage aging in sequence to obtain the high-strength, high-toughness, and heat-resistant cast aluminum alloy; the two-stage solution treatment includes primary solution treatment and secondary solution treatment, the temperature of the primary solution treatment is < the temperature of the secondary solution treatment, the difference being 22~38℃; the two-stage aging treatment includes primary aging and secondary aging, the temperature of the primary aging treatment is < the temperature of the secondary aging treatment, the difference being 40~65℃.

[0011] Further, in step S2, the casting machine is a low-pressure casting machine or a differential pressure casting machine; and / or the casting temperature is 680~700℃; and / or step S2 includes: placing the aluminum alloy melt in the holding furnace of the casting machine, connecting the aluminum alloy melt to the mold through a riser pipe, introducing dry compressed air to form pressure on the surface of the aluminum alloy melt, so that the aluminum alloy melt is sequentially raised, filled, pressurized, held and depressurized along the riser pipe from bottom to top to obtain a casting blank.

[0012] Further, in step S2, the diameter of the riser pipe of the casting machine is 90~160mm; and / or the pressure of low-pressure casting is 110~160KPa, the riser rate is 1.2~1.7KPa / s, the filling rate is 0.8~1.1KPa / s, the crusting pressure increase value is 10~20KPa, the crusting time is 10~20s, and the holding time is 200~800s; and / or the pressure of the lower tank of differential pressure casting is 400~500KPa, the pressure of the upper tank is 300~400KPa, the pressure difference is 100~200KPa, the riser rate is 1.1~1.6KPa / s, the filling rate is 0.9~1.1KPa / s, the crusting pressure increase value is 12~18KPa, the crusting time is 10~20s, and the holding time is 300~600s.

[0013] Further, in step S3, the temperature of the first solution is 500~520℃, and the holding time is 4~8h; and / or the temperature of the second solution is 535~545℃, and the holding time is 8~12h; and / or the quenching is water quenching, the temperature is 60~80℃, and the quenching transfer time is 5~10s.

[0014] Furthermore, in step S3, the temperature for the first aging is 100~120℃ and the holding time is 4~6h; and / or the temperature for the second aging is 150~170℃ and the holding time is 10~14h.

[0015] Further, step S1 includes: step S11, preparing materials according to the composition of high-strength, toughness and heat-resistant cast aluminum alloy to obtain aluminum ingot, AlCu master alloy, AlMn master alloy, Mg ingot, AlZr master alloy, AlV master alloy, AlTi master alloy, AlB master alloy, AlCd master alloy, AlAg master alloy and AlRE master alloy; step S12, holding the aluminum ingot at the first holding temperature, and sequentially adding AlCu master alloy, AlMn master alloy, AlV master alloy and AlCd master alloy at the first holding temperature to obtain a first intermediate melt; step S13, holding the first intermediate melt at the second holding temperature, and sequentially adding AlZr master alloy, AlAg master alloy, Mg ingot and AlRE master alloy at the second holding temperature, and after complete melting, adding AlTi master alloy and AlB master alloy to obtain a second intermediate melt; step S14, passing argon gas through the second intermediate melt for refining, adding a refining agent to obtain an aluminum alloy melt.

[0016] Furthermore, the temperature of the first heat preservation is 720~740℃; and / or the temperature of the second heat preservation is 700~720℃; and / or the temperature of refining is 710~730℃, the time is 15~25min, and the rotation speed is 500~700r / min.

[0017] By employing the technical solution of this invention, and using a multi-element microalloying composition of Cu, Mn, Mg, Zr, V, Ti, B, Cd, Ag, RE, etc., the resulting high-strength, high-toughness, and heat-resistant cast aluminum alloy exhibits significantly improved tensile strength at both room temperature and high temperature, more moderate elongation, and simultaneously improved yield strength. The grain size is also significantly refined, achieving a balance of high strength, high toughness, and heat resistance, thus meeting the requirements of aerospace, defense, and military industries for high-performance aluminum alloy castings. Attached Figure Description

[0018] 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:

[0019] Figure 1 The electron backscatter diffraction (EBSD) microstructure of Embodiment 1 of the present invention in the T6 state is shown;

[0020] Figure 2 The image shown is a scanning electron microscope (SEM) image of Embodiment 1 of the present invention in the T6 state;

[0021] Figure 3 The image shown is a transmission electron microscope (TEM) image of Embodiment 1 of the present invention in the T6 state;

[0022] Figure 4SEM images of the ZL205A alloy in the T6 condition of Comparative Example 1 are shown. Detailed Implementation

[0023] 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.

[0024] As described in the background section of this invention, existing technologies suffer from the problem of aluminum alloys failing to simultaneously achieve high strength, high toughness, excellent high-temperature performance, and good casting properties. To address this issue, in a typical embodiment of this invention, a high-strength, high-toughness, and heat-resistant castable aluminum alloy is provided, comprising, by weight percentage: Cu 4.5-5.5%, Mn 0.4-0.8%, Mg 0.15-0.3%, Zr 0.15-0.25%, V 0.08-0.15%, Ti 0.08-0.15%, B 0.005-0.02%, Cd 0.15-0.25%, Ag 0.05-0.20%, RE 0.05-0.20%, Fe 0.01-0.05%, Si 0.01-0.03%, with the balance being Al and unavoidable impurity elements, where a single impurity element ≤0.03% and the total impurity element ≤0.12%. As an example, common aluminum alloy impurities include C, P, and S.

[0025] Cu, as a major alloying element and core strengthening element, has its content controlled within the aforementioned range to balance strengthening effect, segregation tendency, casting performance, and synergistic effect with other elements. Cu forms a high-density, dispersed θ′ metastable phase in the aluminum matrix. This phase can precipitate uniformly during aging, producing a strong precipitation strengthening effect. When the Cu content is below 4.5 wt.%, the volume fraction of the precipitated phase is insufficient, and the strengthening effect is not obvious; when the Cu content is above 5.5 wt.%, coarse equilibrium phase θ is easily formed, which reduces the strengthening efficiency and impairs plasticity. This invention controls Cu within the aforementioned range, obtaining sufficient driving force for θ′ phase precipitation while avoiding uneven decomposition of supersaturated solid solutions, thereby achieving high levels of room temperature tensile strength and 200°C high-temperature tensile strength in the casting.

[0026] Furthermore, although increasing the Cu content will reduce the elongation to some extent, by strictly controlling the upper limit of the Cu content and adding elements such as Mn, Zr, Ti, and Cd, it is still possible to effectively refine the grains and change the distribution morphology of the precipitated phases. Mn can partially form Al-Cu-Mn ternary intermetallic compounds with Cu (such as Al...). 20The Al-Cu-Mn3 phase exhibits excellent thermal stability, pinning grain boundaries and subgrain boundaries at high temperatures, significantly inhibiting grain growth and dislocation movement. Furthermore, compared to the single Al-Cu phase, the Al-Cu-Mn phase has a lower coarsening rate, thus enabling the alloy to possess good heat resistance. Cd promotes the uniform dispersion and precipitation of the θ′ phase and delays its coarsening during aging; Zr and Ti can form coherent Al3Zr and Al3Ti phases, further strengthening the matrix. This alloy composition allows the alloy to achieve an excellent balance between high strength and good plasticity while maintaining high strength. This invention controls Cu within the aforementioned range, fully utilizing the precipitation strengthening effect of the Al2Cu phase, while synergistically incorporating Mn, Zr, Ti, and Cd elements to overcome the shortcomings of poor plasticity and severe segregation in high-copper alloys, achieving a synergistic improvement in high strength, toughness, and heat resistance.

[0027] Mn, as an important auxiliary alloying element, has its content controlled within the aforementioned range to fully balance grain refinement, heat-resistant phase formation, plasticity retention, and synergistic effects with other elements. In aluminum alloys, Mn can form dispersed Al6Mn phases and react with copper to form Al-Cu-Mn ternary intermetallic compounds. These phases act as heterogeneous nucleation sites or pinned at grain boundaries during solidification, effectively inhibiting grain growth. Controlling the Mn content within the aforementioned range refines the grain size of the casting body and significantly reduces or even eliminates coarse casting columnar crystals, forming a uniform equiaxed grain structure. Compared to Al-Cu alloys with no Mn or low Mn content, the grain refinement effect significantly improves the alloy's strength and toughness. Furthermore, during solution treatment and high-temperature service, the dispersed Al6Mn and Al-Cu-Mn phases effectively pin dislocations and subgrain boundaries, significantly inhibiting the growth of recrystallized grains. This is beneficial for improving the alloy's structural stability at 200℃ and avoids a sharp decrease in strength due to recrystallization.

[0028] Furthermore, excessively high Mn content can reduce elongation due to the formation of coarse manganese-containing phases. Therefore, strictly controlling Mn within the aforementioned range is crucial to avoid excessive embrittlement. At this content, the fine-grained strengthening and dispersion strengthening contributions from Mn are significant, resulting in high room temperature elongation and 200°C elongation, achieving a balance between high strength and high plasticity. This invention controls Mn within the aforementioned range and, in conjunction with elements such as Zr and V, forms a multi-layered dispersion strengthening network, allowing the alloy to maintain a fine-grained structure even at high temperatures. Mn can also form the Al6(Fe,Mn) phase with iron impurities, transforming the acicular β-AlFeSi phase into a blocky or Chinese character-shaped structure, thereby eliminating the brittleness hazard of iron impurities. By controlling the Fe content within a low range, the fracture toughness and fatigue resistance of the alloy are further improved. The above-mentioned composition of the present invention is beneficial to give full play to the multiple functions of Mn in refining grains, inhibiting recrystallization, forming a high heat-resistant Al-Cu-Mn phase, reducing segregation and improving the morphology of impurities, and together with Cu, Zr, Ti, Cd and other elements, it forms the microstructure basis of a high-strength, high-toughness and heat-resistant aluminum alloy.

[0029] As a key synergistic alloying element, Mg's content is controlled within the aforementioned range to fully balance age-hardening effects, grain boundary structure regulation, high-temperature stability, and resistance to stress corrosion. Firstly, Mg combines with copper in the aluminum matrix to form the S phase (Al₂CuMg). This phase co-precipitates with the θ′ (Al₂Cu) phase, producing a composite strengthening effect. Compared to the single θ′ phase, the S phase exhibits higher thermal stability and resistance to coarsening. Controlling Mg within the aforementioned range induces sufficient S phase nucleation while avoiding excessive Mg leading to continuous S phase precipitation at grain boundaries, which would compromise toughness. Simultaneously, an appropriate amount of Mg lowers the critical nucleation energy of the θ′ phase, promoting its uniform and fine precipitation, and reducing the width of the precipitation-free zone near the grain boundaries. The narrowed precipitation-free zone (PFZ) effectively suppresses grain boundary slip and crack initiation, allowing the alloy to maintain high strength while achieving higher room-temperature elongation, thus achieving a balance between precipitation strengthening and grain boundary toughness.

[0030] Meanwhile, the addition of Mg increases the coarsening activation energy of the S and θ′ phases, delaying the growth and dissolution of precipitates during high-temperature service. Furthermore, a small amount of Mg synergistically with Cd further promotes the uniform dispersion of the θ′ and S phases, inhibits the formation of coarse equilibrium phases at grain boundaries, and, combined with the grain-refining effects of Zr and Ti, achieves a fine, discontinuous distribution of precipitates at the alloy grain boundaries, thereby simultaneously improving strength, plasticity, and fatigue resistance. The above-mentioned composition of this invention is beneficial for leveraging the combined strengthening of the S and θ′ phases, refining precipitates, narrowing the PFZ, and improving high-temperature stability, while avoiding the stress corrosion cracking and plasticity reduction problems that are common in high-Mg alloys, achieving a synergistic improvement in strength, plasticity, heat resistance, and corrosion resistance.

[0031] As a key microalloying element, Zr's content is controlled within the aforementioned range to fully leverage its grain refinement, recrystallization inhibition, high-temperature strengthening, and synergistic effects with other elements. By controlling Zr within this range to place it in a supersaturated solid solution region, the dispersed submicron-sized Al3Zr particles, possessing an L12 ordered structure, are coherent or semi-coherent with the aluminum matrix during solidification and subsequent heat treatment. These particles can act as heterogeneous nucleation cores, significantly refining the grains. Simultaneously, Al3Zr particles exhibit extremely high thermal stability, effectively pinning dislocations and subgrain boundaries, inhibiting recrystallization grain growth during solution treatment and high-temperature service, and preventing a sharp drop in strength due to recrystallization.

[0032] Furthermore, Zr atoms can segregate at the interface between the θ′ phase and the matrix, reducing the interfacial energy and thus suppressing the coarsening and dissolution of the θ′ phase during high-temperature aging and service. Simultaneously, Zr precipitates together with Mn and V elements in the alloy, forming various thermally stable dispersed phases such as Al3Zr, Al6Mn, and Al3V. These phases constitute a three-dimensional pinning network within the grains and at grain boundaries, which can suppress high-temperature creep deformation at multiple scales. The above-mentioned composition of this invention is beneficial for fully utilizing the multiple effects of Zr in refining grains, suppressing recrystallization, delaying θ′ phase coarsening, and improving fatigue and creep resistance, forming the microstructure basis of a high-strength, high-toughness, and heat-resistant aluminum alloy together with Cu, Mn, Ti, V, and other elements.

[0033] As a key microalloying element for improving high-temperature stability and weldability, the content of v is controlled within the aforementioned range to fully leverage its role in forming a highly stable Al3V dispersed phase, pinning grain boundaries, inhibiting recrystallization, and improving the microstructure of the weld heat-affected zone. V has low solid solubility in the aluminum matrix, and during solidification and subsequent heat treatment, it can precipitate submicron-sized Al3V dispersed phases. These phases possess extremely high thermal stability and resistance to coarsening, effectively pinning grain boundaries and subgrain boundaries. V can refine as-cast grains and improve the distribution of grain boundary precipitates, resulting in the formation of fine equiaxed grain structures during weld pool solidification, reducing the continuous distribution of coarse columnar grains and low-melting-point liquid films, thereby significantly reducing the tendency for weld hot cracking. Simultaneously, the Al3V dispersed phase can effectively hinder dislocation climb and grain boundary slip at high temperatures, exhibiting a significant dispersion strengthening effect.

[0034] Furthermore, controlling the V content to a low range can prevent excessive V from forming coarse primary phases that impair plasticity. The above-mentioned composition of the present invention is beneficial for fully utilizing the functions of V in forming a highly stable Al3V dispersed phase, inhibiting grain boundary migration, constructing a multi-dispersion strengthening network, improving weldability, and extending high-temperature service life.

[0035] Ti, as a highly efficient grain refiner and auxiliary strengthening element, is controlled within the aforementioned range to form TiB2 particles with B. This fully leverages Ti's heterogeneous nucleation effect, anti-attenuation properties, and synergistic refining effect with Zr, thereby improving mechanical properties. Ti and B react in the melt to generate high-density, dispersed TiB2 particles. TiB2 exhibits excellent lattice matching with the aluminum matrix, and its nucleation efficiency is far higher than that of the Al3Ti phase formed by adding Ti alone. During solidification, these particles act as heterogeneous nucleation sites for α-Al, significantly refining the as-cast grains and eliminating columnar crystals to form a fully equiaxed crystal structure. When Ti is used alone, Al3Ti particles tend to aggregate and grow or settle as the melt holding time increases, leading to a decrease in the refining effect. By introducing B to form TiB2, the TiB2 particles exhibit extremely high thermal stability, maintaining a fine and dispersed state even after prolonged holding at suitable melting temperatures, without significant agglomeration or precipitation. This ensures that the alloy of this invention achieves consistent grain refining effects in large-scale continuous production, resulting in a wider process window.

[0036] Meanwhile, the addition of a small amount of Ti can alter the precipitation behavior of θ′ (Al2Cu) and S phase (Al2CuMg), promoting a more dispersed and uniform distribution of precipitates. Furthermore, through grain refinement, it further shortens the width of the grain boundary non-precipitate zone and inhibits the continuous precipitation of coarse equilibrium phases at grain boundaries, thus maintaining good fracture toughness under high strength. Moreover, the fine equiaxed grain structure makes dendrite overlap more uniform in the later stages of solidification, resulting in a continuous liquid film distribution and significantly reducing the alloy's susceptibility to hot cracking. This invention, by controlling Ti within the aforementioned range and synergistically proportioning it with B and Zr, effectively leverages the heterogeneous nucleation of TiB2, resistance to grain refinement degradation, synergistic effects with Zr, grain refinement and toughening, improvement of the weld heat-affected zone, and reduction of hot cracking tendency.

[0037] Bo (B), as a highly efficient grain-refining auxiliary element, is controlled within the aforementioned range to synergistically form TiB2 particles with Ti, avoiding the coarse, B-rich phase that easily forms when B is added alone. This facilitates the full utilization of B's ​​heterogeneous nucleation effect, prevents microstructure coarsening, and improves the synergistic refining effect with Ti. The fine, dispersed TiB2 particles formed by B and Ti in the melt can act as heterogeneous nucleation sites for aluminum grains during solidification, eliminating coarse columnar crystals and forming a uniform equiaxed crystal structure. When the B content is too high, excess B will form coarse AlB2 phases with Al or Ti agglomerates with Ti, accumulating and distributing in the matrix, becoming stress concentration sources and crack initiation points, leading to a decrease in compressive strength and a significant reduction in elongation. B promotes the formation of fine equiaxed crystals through the heterogeneous nucleation effect of TiB2, resulting in more uniform dendrite overlap in the later stages of solidification, continuous liquid film distribution, and a significant reduction in the tendency for hot cracking and shrinkage porosity. This invention controls B within the above-mentioned range and synergistically forms TiB2 with Ti, which can fully leverage the excellent effects of B, such as efficient heterogeneous nucleation, resistance to refinement decay, avoidance of coarse phase formation, synergistic multi-refinement, and improved strength and toughness.

[0038] Ag, as an aging acceleration and precipitation phase control element, is controlled within the aforementioned range to fully balance accelerating aging response, refining θ′ phase size, improving grain boundary segregation, and enhancing high-temperature stability. In the aluminum matrix, Ag can lower the critical nucleation energy of the θ′ phase, providing an additional heterogeneous nucleation interface, thereby significantly accelerating aging precipitation kinetics. This reduces heat treatment costs, increases production efficiency, and avoids grain coarsening that may result from prolonged aging. Ag atoms can segregate at the interface between the θ′ phase and the matrix, reducing interfacial energy and inhibiting rapid θ′ phase growth, resulting in a finer, more uniform, and denser precipitate size distribution. The smaller θ′ phase produces a stronger precipitation strengthening effect. Simultaneously, Ag can alter the distribution behavior of solute atoms such as Cu and Mg at grain boundaries, inhibiting the continuous precipitation of coarse equilibrium phases at grain boundaries and promoting intragranular precipitation. This significantly reduces the PFZ width; the narrower and discontinuous PFZ effectively inhibits grain boundary crack initiation, achieving a balance between strength and toughness. In addition, Ag can reduce the anisotropy of mechanical properties caused by texture or directional precipitation by promoting the uniform distribution of precipitates on multiple slip systems. This is especially important for castings with complex shapes and subjected to multi-directional loads, and can significantly improve the reliability and consistency of the design.

[0039] Furthermore, the combined action of Ag and Cd can further promote the uniform nucleation of the θ′ phase and inhibit its coarsening; when coexisting with Mg, Ag can promote the uniform distribution of the S phase; the synergistic effect of multiple elements is beneficial to maximizing the precipitation strengthening effect, while avoiding the adverse effects of excessive single elements. The above composition of the present invention is beneficial to fully utilize the effects of Ag in accelerating aging response, refining the θ′ phase, reducing grain boundary segregation, improving anisotropy, enhancing high-temperature stability and corrosion resistance, and together with Cu, Mg, Cd, Zr and other elements, forms a microalloying system for high-strength, high-toughness and heat-resistant aluminum alloys.

[0040] Rare earth elements (REs) can refine grains, regulate impurity phases, and improve damping performance and corrosion resistance. Taking REs including La and Ce as an example, La and Ce have extremely low solid solubility in aluminum matrices. During solidification, they precipitate diffusely as La-rich and Ce-rich phases, respectively. When both coexist, the resulting (La,Ce) composite intermetallic compound exhibits higher thermal stability and better lattice matching, serving as highly efficient heterogeneous nucleation sites. RE atoms can co-segregate at grain boundaries, reducing grain boundary free energy and hindering the excessive segregation of solute atoms such as Cu and Mg at grain boundaries, thereby suppressing the precipitation of continuous, coarse θ or S phases at grain boundaries. Furthermore, the rare earth-rich phase can pin dislocations and grain boundaries at high temperatures, assisting the dispersion strengthening effect of elements such as Zr, Mn, and V. Rare earth elements have a strong affinity for impurities such as oxygen, hydrogen, and sulfur in the melt, forming high-melting-point compounds that float or sink, thus deeply purifying the melt. Rare earth elements can also reduce the surface tension of the aluminum melt, improving its filling ability. The above-mentioned composition of this invention is beneficial for fully utilizing the RE composite to form an efficient heterogeneous nucleation core, synergistically regulate harmful Fe phases, suppress grain boundary segregation, optimize damping performance, improve corrosion resistance, and improve casting fluidity. Together with elements such as Cu, Mn, Mg, Zr, V, Ti, B, Cd, and Ag, it forms a microalloying system for high-strength, high-toughness, and heat-resistant aluminum alloys.

[0041] The high-strength, high-toughness, and heat-resistant cast aluminum alloy of this invention has a specific elemental composition, resulting in significantly improved performance. Firstly, it exhibits excellent room-temperature mechanical properties: the room-temperature tensile strength of the casting body reaches 490~520MPa, the yield strength is 400~420MPa, and the elongation is 8~12%. Compared to traditional Al-Cu cast aluminum alloys, this invention significantly improves plasticity while maintaining high strength, achieving a good balance between high strength and high toughness. Secondly, it possesses outstanding high-temperature (200℃) tensile properties: under 200℃ testing conditions, the high-temperature tensile strength of the casting body reaches 350~390MPa, and the elongation is 12.5~16%. This indicates that the alloy can still maintain high load-bearing capacity and good plasticity at around 200℃, making it suitable for applications with stringent heat resistance requirements, such as aerospace engine accessories, automotive turbocharger housings, and high-temperature structural components. Thirdly, it possesses a fine-grained microstructure: the average grain size of the casting body is controlled at 60~80μm, far smaller than that of conventional cast aluminum alloys (typically 100~200μm). This fine-grained microstructure helps to simultaneously improve strength and plasticity, and enhances the alloy's fatigue resistance and resistance to hot cracking. Furthermore, by controlling the Fe content to 0.01~0.05% and the Si content to 0.01~0.03%, this invention effectively avoids the formation of brittle acicular AlFeSi phases, significantly improving the alloy's fracture toughness and fatigue resistance.

[0042] In a preferred embodiment, RE comprises La and Ce, and the weight ratio of La to Ce is 1:(1~3); and / or the weight ratio of Mn to Cu is 0.08~0.15; and / or the weight ratio of Zr to Ti is 1.2~2.0; and / or the weight ratio of Ti to B is 6~26.

[0043] La has a strong affinity for Fe, preferentially forming stable blocky or spherical Al–La–Fe phases, thus consuming Fe atoms and preventing Fe from forming brittle Cu2FeAl7 phases with Cu and Al. Ce can further alter the morphology of the remaining Fe phase, transforming it into a Chinese character-shaped or granular form. The synergistic effect of La and Ce minimizes the harmful effects of the iron phase. Compared to adding La or Ce alone, the simultaneous addition of La and Ce, with the weight ratio controlled within the aforementioned range, results in the highest nucleation density and a finer average grain size in the casting. Furthermore, the composite addition of La / Ce can more effectively narrow the PFZ width than single rare earth elements and cause discontinuous distribution of grain boundary precipitates, significantly improving the fracture toughness and fatigue resistance of the alloy. Controlling the La / Ce weight ratio within the aforementioned range is beneficial for controlling the Ce proportion, making it less likely for the Ce-rich phase to aggregate and grow, thus significantly improving damping performance; simultaneously controlling the La proportion further improves the grain refinement effect and corrosion resistance.

[0044] Maintaining the Mn / Cu weight ratio within the above range is preferable, as it promotes the partial formation of Al-Cu-Mn ternary intermetallic compounds (such as Al...). 20 (Cu2Mn3 phase). Furthermore, the Mn / Cu weight ratio was strictly controlled at 0.08–0.15 to ensure that the stoichiometric ratio of Mn to Cu was within the optimal range for forming the Al-Cu-Mn ternary phase. This ternary phase possesses a higher coarsening activation energy than the pure Al2Cu phase, and exhibits an extremely low coarsening rate under long-term thermal exposure at 200℃. Through the high-temperature pinning effect of this phase, the alloy's tensile strength at 200℃ is consistently higher. In addition, an appropriate amount of Mn can alter the partition coefficient of Cu at the solid-liquid interface, reducing the degree of copper dendrite segregation and resulting in a more uniform distribution of the Al2Cu precipitate phase.

[0045] Ti primarily provides nucleation sites by forming the Al3Ti phase, but the grain refinement effect of Ti is easily weakened by Zr; while the nucleation efficiency of Al3Zr particles from Zr alone is slightly lower. By optimizing the Zr / Ti ratio, Al3Zr and Al3Ti complement each other. Controlling the Zr / Ti weight ratio within the aforementioned range is ideal. At this ratio, Ti mainly provides high-density nucleation cores as TiB2 or a small amount of Al3Ti, while Zr provides long-term grain boundary pinning and recrystallization inhibition as the dispersed Al3Zr phase. The two do not weaken each other (excess Ti inhibits the precipitation of pinning phases in Zr, and excess Zr occupies nucleation sites and inhibits the nucleation activity of Ti), and they achieve synergy between nucleation and pinning (Al3Ti provides high-density effective nucleation cores, while Al3Zr provides long-term pinning). This synergy maximizes the grain refinement effect, ensuring the alloy maintains a fine and uniform grain structure in the as-cast, solution-treated, and aged states.

[0046] It is advisable to control the weight ratio of Ti to B within the above range, which is beneficial to optimize the generation efficiency and nucleation density of TiB2 particles, give full play to the heterogeneous nucleation effect of TiB2, further refine the grains, form a uniform equiaxed structure, eliminate columnar crystals, thereby improving strength and plasticity and reducing the tendency of hot cracking.

[0047] To further improve the overall performance of aluminum alloys, in a preferred embodiment, the high-strength, high-toughness, and heat-resistant cast aluminum alloy comprises, by weight percentage: Cu 4.8~5.2%, Mn 0.5~0.7%, Mg 0.2~0.28%, Zr 0.18~0.24%, V 0.10~0.14%, Ti 0.09~0.14%, B 0.006~0.015%, Cd 0.18~0.24%, Ag 0.08~0.15%, RE 0.08~0.18%, Fe 0.01~0.05%, Si 0.01~0.03%, with the balance being Al and unavoidable impurity elements.

[0048] Existing aluminum alloys also suffer from poor compatibility with manufacturing processes. Most existing smelting, refining, casting, and heat treatment processes are based on conventional Al-Si or Al-Cu alloys, which are difficult to meet the requirements of new multi-element microalloyed Al-Cu alloys for melt purity, filling pressure, temperature control, and precipitation kinetics. This results in castings with many internal defects, coarse microstructure, and unstable performance. Therefore, in another typical embodiment of the present invention, a method for preparing the above-mentioned high-strength, high-toughness, and heat-resistant cast aluminum alloy is also provided, comprising the following steps: Step S1, batching, melting, and refining the high-strength, high-toughness, and heat-resistant cast aluminum alloy in sequence to obtain an aluminum alloy melt; Step S2, transferring the aluminum alloy melt to a casting machine for casting, including low-pressure casting or differential pressure casting, to obtain a cast blank; Step S3, subjecting the cast blank to two-stage solution treatment, quenching, and two-stage aging in sequence to obtain a high-strength, high-toughness, and heat-resistant cast aluminum alloy; the two-stage solution treatment includes primary solution treatment and secondary solution treatment, the temperature of the primary solution treatment is < the temperature of the secondary solution treatment, with a difference of 22~38℃; the two-stage aging treatment includes primary aging and secondary aging, the temperature of the primary aging treatment is < the temperature of the secondary aging treatment, with a difference of 40~65℃.

[0049] Aluminum alloy melt with a specific composition is transferred to a holding furnace in a casting machine for low-pressure or differential-pressure casting to obtain a cast blank. With Cu content within the range specified in this invention, combined with specific smelting and casting processes, the micro-segregation of copper can be controlled at a low level, avoiding the formation of a continuous brittle network θ phase at grain boundaries. Improved microstructure uniformity not only benefits the consistency of mechanical properties but also reduces the tendency for hot cracking, thus improving casting performance. Low-differential-pressure casting can effectively reduce casting defects such as porosity and shrinkage, ensuring the consistency of casting density and mechanical properties.

[0050] The cast blank is subjected to vibration-induced sand removal, and the riser and gating system are sawn off. Then, shot blasting is performed to remove adhering sand and burrs, preventing stress concentration. The blank is then milled to the predetermined specifications to ensure dimensional accuracy. Next, heat treatment is performed. Because the alloy of this invention contains both low-melting-point and high-melting-point phases, and multiple non-equilibrium eutectic phases exist simultaneously, directly heating to a high temperature for solution treatment would cause the low-melting-point phases to melt first, forming a localized "overheated" structure, severely impairing mechanical properties. Therefore, the heat treatment employs a two-stage solution treatment and a two-stage aging process.

[0051] This invention achieves the effects of safe dissolution of low-melting-point phases, thorough dissolution of high-melting-point phases, low residual stress, and good matching with multi-component microalloying systems through a two-stage solution treatment process. The first solution treatment dissolves the low-melting-point phases, while the second treatment further promotes the dissolution and homogenization of the insoluble phases. Quenching preserves high supersaturation. Since the low-melting-point phases have already been pre-dissolved in the first treatment, there is no risk of overheating during the second treatment. After this treatment, subsequent aging yields higher density θ′ and S phase precipitation. Through the two-stage solution treatment and quenching, the alloy acquires a high density of vacancies (quenching vacancies), providing nucleation sites for subsequent aging. The solute atoms are evenly distributed, resulting in fine grains and clean grain boundaries. This microstructure promotes high-density, uniform nucleation of θ′ and S phases during subsequent secondary aging, while suppressing the widening of the non-precipitate zone at grain boundaries and the precipitation of coarse equilibrium phases. Ultimately, the alloy simultaneously achieves high strength, high plasticity, and excellent high-temperature performance.

[0052] In addition, this invention adopts the method of continuing to heat to the secondary solution temperature after the first solution is completed, instead of cooling after being taken out of the furnace and then reheating. This avoids the precipitation of intermediate phases during the cooling process after the first solution, which would require re-dissolving during the secondary solution, wasting time and energy. It also reduces the number of times the casting is exposed outside the furnace and quenched and transferred, reducing the risk of oxidation and deformation. The continuous heating process makes the microstructure evolution more continuous and controllable.

[0053] A two-stage aging process enables separate control of nucleation and growth. The first aging promotes low-temperature nucleation, while the second aging promotes high-temperature growth. Single-stage high-temperature aging results in fewer nucleation sites, sparse and coarse precipitates, lower strength, wider grain boundary free zone (PFZ), and poor plasticity. Conversely, single-stage low-temperature aging produces excessively fine precipitates, leading to insufficient strengthening. This is because single-stage high-temperature aging easily causes solute atoms at grain boundaries to diffuse into the grain, forming wide PFZs that become preferential pathways for crack propagation. In contrast, with two-stage aging, the first low-temperature nucleation pre-forms a large number of fine precipitates near the grain boundaries, consuming the solute in these areas. The second aging process then limits the PFZ width to a narrower range. This narrow PFZ, combined with the discontinuous distribution of equilibrium phases at grain boundaries, maintains a high room-temperature elongation of the alloy and improves its resistance to fatigue crack propagation. This invention achieves separation and control of low-temperature high-density nucleation and high-temperature controllable growth by limiting a two-stage aging process, giving full play to the synergistic strengthening effect of Ag and Cd, narrowing the grain boundary precipitation-free zone, ensuring the performance consistency of complex castings, and avoiding the decrease in plasticity caused by over-aging.

[0054] On the one hand, the segregation of Ag and Cd atoms in the first aging stage promotes the uniform nucleation of high-density GP zones and θ″ precursors. In the second aging stage, the size of the fine θ′, S, and Ω phases is controlled within the optimal strengthening range, resulting in high-density, diffusely distributed precipitates without significant coarsening. Simultaneously, the width of the non-precipitation zone at grain boundaries is significantly narrowed to ≤50nm, suppressing the precipitation of continuous coarse equilibrium phases at grain boundaries. This allows the alloy to achieve the highest precipitation strengthening increment while effectively avoiding grain boundary embrittlement and over-aging softening. On the other hand, the progressive precipitation process of the two-stage aging process helps improve the uniformity of the aging response in various parts of complex castings, avoiding performance fluctuations caused by uneven temperature fields in single-stage aging. Ultimately, the alloy achieves high levels of room-temperature tensile strength, yield strength, elongation, and high-temperature mechanical properties, realizing a match between high strength, high toughness, and excellent high-temperature performance, and improving the consistency and reliability of mechanical properties between batches.

[0055] The multi-stage heat treatment process can promote the uniform precipitation of θ′(Al2Cu) phase, Al3Zr, Al3Ti, Al6Mn, V-containing and rare earth-containing dispersed phases, and suppress the continuous distribution of coarse equilibrium phases at grain boundaries. By utilizing the good synergistic effect of precipitation strengthening and grain boundary strengthening, fine, dispersed, and thermally stable θ′ and S phases are generated. At the same time, the widening of the non-precipitation zone at grain boundaries and the precipitation of coarse equilibrium phases are suppressed, which significantly improves the room temperature and high temperature performance of the casting body, and obtains the finished casting, namely the high strength, toughness and heat resistance cast aluminum alloy of this invention.

[0056] Furthermore, the specific Cu content allows the alloy to withstand multi-stage solution aging treatment. A higher Cu content ensures the formation of a supersaturated solid solution after secondary solution treatment, while a moderate Cu content prevents localized melting during the solution process. After multi-stage solution aging treatment, the θ′ phase exhibits high density, uniform distribution, and good thermal stability, resulting in optimal strength, plasticity, and heat resistance properties. The Mg content, combined with multi-stage aging treatment, ensures the alloy maintains high high-temperature tensile strength even after long-term exposure at high temperatures. After preparation, non-destructive testing (X-ray or ultrasonic testing) and welding repair processes can be performed to ensure the casting is defect-free and achieves a high yield.

[0057] This invention optimizes the characteristics of the Al-Cu multi-component microalloying system. By combining the synergistic effect of multi-component microalloying with optimized low differential pressure casting and multi-stage heat treatment, a high-strength, high-toughness, and heat-resistant cast aluminum alloy is obtained. The mechanical properties of the alloy body at room temperature and 200℃ are significantly improved, and the grain size is refined, meeting the needs of aerospace, defense and military industries for high-strength, high-toughness and heat-resistant aluminum alloy castings.

[0058] To further reduce alloy solidification shrinkage and porosity, in a preferred embodiment, in step S2, the casting machine is a low-pressure casting machine or a differential pressure casting machine; and / or the casting temperature is 680~700℃; and / or step S2 includes: placing the aluminum alloy melt in the holding furnace of the casting machine, connecting the aluminum alloy melt to the mold through a riser pipe, introducing dry compressed air to form pressure on the surface of the aluminum alloy melt, so that the aluminum alloy melt sequentially rises, fills the mold, pressurizes, holds pressure and depressurizes along the riser pipe from bottom to top to obtain a cast blank.

[0059] To address the characteristics of Al-Cu alloys, such as a wide solidification temperature range, high tendency to hot cracking, and susceptibility to shrinkage porosity and segregation, this invention preferably employs a nonlinear pressurization process. Low differential pressure casting allows the molten metal to smoothly fill the mold from bottom to top under pressure, reducing turbulent air entrapment. During the filling stage, the rate at which the molten metal enters the mold cavity needs to be adjusted according to the casting wall thickness and complexity. Immediately after filling, a shell-forming pressurization process is performed and maintained (the holding time is the shell-forming time), causing the casting surface to rapidly solidify under pressure, forming a dense solidified shell. The molten metal forms a temperature gradient within the mold cavity, from thin-walled to thick-walled, through the gating, and then to the riser pipe, achieving sequential solidification.

[0060] In a preferred embodiment, in step S2, the diameter of the riser pipe of the casting machine is 90~160mm; and / or the pressure of low-pressure casting is 110~160KPa, the riser rate is 1.2~1.7KPa / s, the filling rate is 0.8~1.1KPa / s, the crusting pressure increase value is 10~20KPa, the crusting time is 10~20s, and the holding time is 200~800s; and / or the pressure of the lower tank of differential pressure casting is 400~500KPa, the pressure of the upper tank is 300~400KPa, the pressure difference is 100~200KPa, the riser rate is 1.1~1.6KPa / s, the filling rate is 0.9~1.1KPa / s, the crusting pressure increase value is 12~18KPa, the crusting time is 10~20s, and the holding time is 300~600s.

[0061] The casting pressure should ideally be controlled within the aforementioned range to optimize the supercooling and feeding capacity of the molten metal. During the rising stage, the molten metal should rise steadily from bottom to top to reduce turbulence and splashing. The aforementioned riser pipe diameter and rising rate promote complete filling at thin-walled sections and reduce impact splashing, allowing the molten alloy to rise in a laminar flow state within the riser pipe. This reduces gas entrapment and secondary oxidation, significantly improving the density and mechanical property consistency of the casting. Controlling the filling rate within the aforementioned range also reduces the formation of chilled crystals and coarse dendrites in the Al-Cu alloy during filling, keeping the molten metal front smooth, reducing core erosion, and extending gas venting time. This allows sufficient time for gas in the cavity to escape through the vent plug, reducing porosity.

[0062] Controlling the shell formation parameters within the above range facilitates the rapid formation of a dense shell layer on the casting surface, reducing internal shrinkage porosity and aluminum melt leakage. This ensures a moderate shell thickness to maintain internal feeding channels, achieving both sufficient density and adequate feeding capacity, effectively suppressing surface shrinkage cavities and macroscopic segregation. The above-mentioned holding time allows the casting to complete most of its solidification under pressure, ensuring sufficient feeding under pressure, thereby controlling shrinkage rate, reducing central shrinkage cavities, and improving casting density and mechanical property consistency. Controlling the nonlinear pressurization parameters of low-pressure casting within the above range promotes stable mold filling, formation of a dense shell layer, sequential solidification for sufficient feeding, reduced hot cracking tendency, reduced macroscopic segregation, and improved dimensional accuracy.

[0063] In a preferred embodiment, in step S3, the temperature of the first solution is 500~520℃ and the holding time is 4~8h; and / or the temperature of the second solution is 535~545℃ and the holding time is 8~12h; and / or the quenching is water quenching at a temperature of 60~80℃ and a quenching transfer time of 5~10s.

[0064] The first solution treatment involves a temperature range within which the AlCd phase and some low-melting-point eutectic phases fully dissolve, eliminating the brittle grain boundary network in the as-cast state. The temperature is below the Al2Cu eutectic point to prevent grain boundary melting. Appropriate heat treatment allows for the initial homogenization of solute atoms such as Cu, Mg, Cd, and Ag, laying the compositional foundation for subsequent high-temperature solution treatment. After the first solution treatment, the alloy retains high-melting-point θ-Al2Cu phase, S-Al2CuMg phase, and dispersed phases containing Zr, V, and Mn. The second solution treatment is performed at a temperature higher than the first but still below the alloy's eutectic point. At this temperature, the remaining θ and S phases fully dissolve, allowing strengthening elements such as Cu and Mg to almost completely enter the aluminum matrix, resulting in a highly supersaturated solid solution. Sufficient heat treatment ensures that the core of thick / large castings also achieves solution saturation, while simultaneously allowing some Zr, V, and Mn to dissolve or the dispersed phases to refine.

[0065] After solution treatment, rapid cooling is necessary to retain high supersaturation. Water quenching helps to more effectively suppress the precipitation of the θ′ phase during cooling, reduces the temperature difference between the casting and the quenching medium, significantly reduces residual quenching stress and deformation cracking tendency, and reduces high-density dislocations and microcracks caused by excessively rapid cooling, which is beneficial for maintaining the alloy's plasticity and fatigue resistance. The quenching transfer time refers to the time from opening the furnace door to the casting being fully immersed in the quenching medium. A shorter quenching transfer time can reduce the tendency of the casting to precipitate coarse equilibrium phases when cooling in air, thereby reducing the consumption of solute atoms in the supersaturated solid solution, maintaining the subsequent aging strengthening effect, and allowing the casting to be directly immersed in water at high temperature in a short time without pre-precipitation, keeping the supersaturation of elements such as Cu and Mg at a high level, providing the maximum driving force for subsequent aging.

[0066] After solution quenching, the alloy yields a highly supersaturated solid solution and a large number of quenching vacancies. In a preferred embodiment, in step S3, the temperature of the first aging is 100~120℃, and the holding time is 4~6h; and / or the temperature of the second aging is 150~170℃, and the holding time is 10~14h.

[0067] The low-temperature range selected for the first aging process is sufficient to promote solute atom diffusion and vacancy migration. At this temperature, a high-density GP region and tiny θ″ phase pre-precipitates can be formed, with extremely high number density. Ag and Cd atoms preferentially agglomerate at the GP region / matrix interface, lowering the critical nucleation energy and further increasing the nucleation density. Appropriate holding allows the nucleation process to proceed fully without causing excessive growth, ensuring a sufficient number of precipitates during subsequent heating. This low-temperature first-stage aging provides a large number of uniform and fine nucleation sites for subsequent high-temperature aging, reducing the coarse and sparse precipitation caused by insufficient nucleation sites during direct high-temperature aging.

[0068] After the first aging, a high-density GP region and θ″ precursor exist in the alloy. Raising the temperature to a high-temperature range, the GP region rapidly transforms into θ″ and θ′ phases. Simultaneously, the θ′ phase grows from its original core, eventually forming needle-like or disc-like θ′ phases, producing a strong precipitation strengthening effect. Within this temperature range, Mg and Cu further form the S phase (Al2CuMg), and Ag promotes the precipitation of the Ω phase, forming a composite strengthening network with the θ′ phase. The segregation of Cd atoms inhibits the lateral coarsening of the θ′ phase during growth, maintaining the fine plate-like morphology of the precipitated phase and delaying its transformation to the equilibrium θ phase. Sufficient heat treatment allows the aging to reach its peak state, at which point the precipitated phase has a uniform size distribution and the largest volume fraction. Furthermore, due to the moderate secondary aging temperature, it does not cause a significant transformation of the θ′ phase to the equilibrium θ phase. The precipitated phase maintains a metastable structure even when serving at 200℃, achieving both maximum strength and high elongation. The aforementioned two-stage aging process facilitates the separation and control of nucleation and growth, resulting in a fine precipitate with optimal size, thereby achieving the highest precipitation enhancement increment.

[0069] In a preferred embodiment, step S1 includes: step S11, preparing materials according to the composition of a high-strength, high-toughness, and heat-resistant cast aluminum alloy to obtain aluminum ingots, AlCu master alloy, AlMn master alloy, Mg ingots, AlZr master alloy, AlV master alloy, AlTi master alloy, AlB master alloy, AlCd master alloy, AlAg master alloy, and AlRE master alloy; step S12, subjecting the aluminum ingots to a first heat preservation temperature, and sequentially adding AlCu master alloy, AlMn master alloy, AlV master alloy, and AlCd master alloy at the first heat preservation temperature to obtain a first intermediate melt; step S13, subjecting the first intermediate melt to a second heat preservation temperature, and sequentially adding AlZr master alloy, AlAg master alloy, Mg ingot, and AlRE master alloy at the second heat preservation temperature, and after complete melting, adding AlTi master alloy and AlB master alloy to obtain a second intermediate melt; step S14, passing argon gas through the second intermediate melt for refining, adding a refining agent to obtain an aluminum alloy melt.

[0070] Since excessive Fe content can lead to the formation of needle-like β-AlFeSi phase or Cu2FeAl7 phase, which severely impairs the toughness and fatigue performance of the alloy, high-purity aluminum ingots are first used to control the content of impurities such as Fe and Si to avoid the formation of needle-like brittle phases. High-melting-point or easily oxidized elements (such as V, Zr, Mn, etc.) are introduced in the form of intermediate alloys to promote their rapid and uniform dissolution and reduce burn-off or segregation caused by the direct addition of pure metals.

[0071] High-purity aluminum ingots are melted in a crucible and held at a high temperature for the first time. AlCu master alloy, AlMn master alloy, AlV master alloy, and AlCd master alloy are added in sequence. Since the primary phases formed by these elements (such as Al2Cu, Al6Mn, Al3V, and AlCd) require a long time to homogenize, and Cd is easily volatilized at high temperatures, adding them later would result in greater losses. This order allows Cu, Mn, V, and Cd to dissolve preferentially and initially form a dispersed distribution, without interfering with the subsequently added Zr, Ti, and other elements, which is conducive to forming a uniform intermediate melt.

[0072] The first intermediate melt is then subjected to a second holding temperature. During this second holding temperature, AlZr master alloy, AlAg master alloy, Mg ingot, and AlRE master alloy are added sequentially. These elements require time to dissolve and form stable Al3Zr, AlAg clusters, Al2CuMg, and RE-rich phases. After complete melting, AlTi master alloy and AlB master alloy are added to obtain the second intermediate melt. TiB2, once formed in the melt, exhibits highly efficient nucleation. Adding it too early may lead to agglomeration and deactivation due to prolonged holding. Adding it last ensures fine dispersion of TiB2 and maximizes nucleation efficiency, while preventing Ti and Zr from mutually weakening the refining effect. Finally, argon gas is introduced into the second intermediate melt for refining, and a refining agent is added to obtain the aluminum alloy melt.

[0073] Since the elements introduced by the refining agent have been included in the proportioning calculation, and considering these elements is not a simple matter of addition and subtraction but requires practical engineering application, a "post-furnace addition" strategy is often adopted in practice. This involves adjusting the pre-furnace matrix elements to near the design lower limit, reserving a buffer space for yield fluctuations, and ensuring that the final composition does not touch the upper limit. This provides sufficient margin for actual process window operation, so the addition of the refining agent will not cause the alloy composition to exceed the design range of this application.

[0074] In a preferred embodiment, the first holding temperature is 720~740℃; and / or the second holding temperature is 700~720℃; and / or the refining temperature is 710~730℃, the time is 15~25min, and the rotation speed is 500~700r / 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 second intermediate melt is 0.4~0.8%.

[0075] The first holding temperature is higher than the Al-Cu eutectic temperature but lower than the Al2Cu phase dissolution temperature, which is beneficial for promoting the rapid melting of Cu, Mn, V, and Cd without overheating and burning. The second holding temperature is lower than the first holding temperature, which can further reduce the high-temperature oxidation or burning of Zr, Ag, Mg, and RE. The higher refining speed can generate fine and uniform argon bubbles, significantly improving degassing efficiency and reducing the hydrogen content of the melt; the appropriate refining time can fully degas and remove impurities while reducing the loss of alloying elements caused by over-refining. In addition, the purity of the melt can be improved by degassing refining (argon, 500~700 r / min, 15~25 min) and density equivalent testing, reducing the porosity and shrinkage defects in the casting and improving the casting yield. Optimizing the rapid transfer of the melt during melting and refining helps to shorten the exposure time of the melt in air, reduce secondary gas absorption and oxidation, and provide a high-purity, low-gas-content high-quality melt for subsequent low differential pressure casting.

[0076] In a preferred embodiment, the aluminum ingot is first held at a temperature in a crucible, which may be one or more of a cast iron crucible, a ceramic crucible, or a graphite crucible. The cast iron crucible is coated with a crucible coating before melting at a temperature of 200-400°C. The crucible coating may include one or more of boron nitride, white corundum, and talc, and the proportions can be adjusted as needed. At high temperatures, the molten metal reacts with the Fe in the uncoated cast iron crucible substrate, causing Fe to rapidly dissolve into the melt and increase the Fe content. Therefore, when using a cast iron crucible, a crucible coating needs to be applied before melting to form a dense refractory insulating layer on the surface of the cast iron crucible. This effectively prevents direct contact between the molten metal and the iron substrate, controls the amount of Fe added to the melt, and ensures that the Fe content of the alloy remains within the designed range. The aforementioned coating temperature facilitates the full curing of the binder in the crucible coating, forming a strong, continuous, and erosion-resistant coating. This ensures optimal bonding strength between the coating and the crucible substrate, allowing it to withstand long-term erosion and corrosion from molten metal at 720-740°C. This extends the service life of the cast iron crucible, reducing replacement frequency and smelting costs. Furthermore, the coating formed at this temperature not only isolates iron but also inhibits the reaction between the molten metal and the crucible wall, reducing the tendency of molten aluminum to absorb hydrogen at high temperatures. The smooth coating surface facilitates slag removal, reduces the incorporation of adhering oxide inclusions into the melt, and fills microcracks and pores on the surface of the cast iron crucible.

[0077] 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.

[0078] Example 1

[0079] Step S11: Prepare the smelting raw materials according to the composition of the high-strength, toughness and heat-resistant cast aluminum alloy (see Table 1 and Table 2). The smelting raw materials specifically include high-purity aluminum ingots, AlCu master alloy, AlMn master alloy, AlV master alloy, AlZr master alloy, AlTi master alloy, AlB master alloy, AlCd master alloy, AlRE master alloy, AlAg master alloy and pure Mg ingots.

[0080] Step S12: Apply crucible coating (boron nitride, white corundum, and talc in a mass ratio of 2:6:2) to a cast iron crucible. The temperature of the cast iron crucible is 250°C during coating. Place the high-purity aluminum ingot into the cast iron crucible and then perform smelting and a first holding temperature of 730°C. Under this temperature, add AlCu master alloy, AlMn master alloy, AlV master alloy, and AlCd master alloy in sequence to obtain the first intermediate melt.

[0081] Step S13: The first intermediate melt is subjected to a second heat preservation at a temperature of 710°C. Then, AlZr master alloy, AlAg master alloy, pure Mg ingot and AlRE master alloy are added in sequence. After complete melting, AlTi master alloy and AlB master alloy are added to obtain the second intermediate melt.

[0082] Step S14: The second intermediate melt is refined by argon gas through a degasser at a temperature of 720°C for 16 minutes at a speed of 510 r / min. After refining, slag is removed. Once the composition is qualified, a refining agent (AlTiB refining agent, 0.5%) is added. The density equivalent of the alloy melt is tested. Once it is qualified (density equivalent is 0.2%), the aluminum alloy melt is quickly transferred to a low-pressure casting machine.

[0083] Step S2: The molten aluminum alloy is placed in the holding furnace of the low-pressure casting machine. The mold and the molten aluminum alloy are connected through a riser pipe with a diameter of 90 mm. Low-pressure casting is then carried out at a casting temperature of 700℃. Dry compressed air is introduced to create pressure on the surface of the molten metal, causing the molten aluminum to rise, fill the mold, increase pressure, hold pressure, and release pressure along the riser pipe from bottom to top. Low-pressure casting is a non-linear pressure increase process with a pressure value of 110 kPa, a riser rate of 1.3 kPa / s, a filling rate of 0.9 kPa / s, a crust formation pressure increase value of 10 kPa, a crust formation time of 12 s, and a holding pressure time of 700 s. Finally, a cast blank is obtained.

[0084] Step S3 involves vibrating the casting blank to remove sand, sawing off the risers and gating system, then shot blasting to remove surface burrs, milling to the predetermined dimensions, followed by solution treatment and aging treatment. The solution treatment consists of a single solution treatment and a double solution treatment. The first solution treatment is heated to 500℃ and held for 8 hours, followed by a second solution treatment at 538℃ for 8 hours. Then, water quenching is performed at 60℃ with a quenching transfer time of 6 seconds. The aging treatment consists of a single aging treatment and a double aging treatment. The first aging treatment is performed at 120℃ for 4 hours, followed by a second aging treatment at 160℃ for 12 hours. Finally, non-destructive testing and repair welding are performed to obtain the finished casting, namely, a high-strength, high-toughness, heat-resistant cast aluminum alloy.

[0085] Examples 2 to 10

[0086] The difference from Example 1 is that the alloy composition and preparation process parameters are different, as detailed in Tables 1 to 5.

[0087] Example 11

[0088] The difference from Example 1 is that in step S2, the aluminum alloy melt is placed in the holding furnace of the differential pressure casting machine, and the mold and the aluminum alloy melt are connected through the riser pipe for differential pressure casting. The pressure of the lower tank is 450 kPa, the pressure of the upper tank is 350 kPa, the pressure difference is 100 kPa, the riser rate is 1.4 kPa / s, the filling rate is 1.0 kPa / s, the shelling pressure is 15 kPa, the shelling time is 15 s, the holding time is 500 s, and finally a cast blank is obtained.

[0089] Example 12

[0090] The difference from Example 1 is that in step S2, the aluminum alloy melt is placed in the holding furnace of the differential pressure casting machine, and the mold and the aluminum alloy melt are connected through the riser pipe for differential pressure casting. The pressure of the lower tank is 400 kPa, the pressure of the upper tank is 300 kPa, the pressure difference is 100 kPa, the riser rate is 1.1 kPa / s, the filling rate is 0.9 kPa / s, the shelling pressure is 12 kPa, the shelling time is 10 s, the holding time is 300 s, and finally a cast blank is obtained.

[0091] Example 13

[0092] The difference from Example 1 is that in step S2, the aluminum alloy melt is placed in the holding furnace of the differential pressure casting machine, and the mold and the aluminum alloy melt are connected through the riser pipe for differential pressure casting. The pressure of the lower tank is 500 kPa, the pressure of the upper tank is 300 kPa, the pressure difference is 200 kPa, the riser rate is 1.6 kPa / s, the filling rate is 1.1 kPa / s, the shelling pressure is 18 kPa, the shelling time is 20 s, the holding time is 600 s, and finally a cast blank is obtained.

[0093] Comparative Examples 1 to 4

[0094] The difference from Example 1 is that the alloy composition and preparation process parameters are different, as detailed in Tables 1 to 5.

[0095] Performance testing:

[0096] The materials of the above embodiments and comparative examples were tested as follows, and the results are shown in Table 6.

[0097] Tensile strength: The room temperature tensile test was 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 was 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.

[0098] Microstructure analysis: The finished casting was analyzed using EBSD scanning electron microscopy, magnified 200x along the casting direction. The average grain size was measured using OIM 8.0 analysis software to evaluate the grain size of the casting samples.

[0099] Non-destructive testing: Non-destructive testing of castings was 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.

[0100] Table 1

[0101]

[0102] Table 2

[0103]

[0104] Table 3

[0105]

[0106] Table 4

[0107]

[0108] Table 5

[0109]

[0110] Table 6

[0111]

[0112] Example 1: EBSD histogram in T6 state (see attached image) Figure 1 The image shows fine equiaxed crystals; SEM image of Example 1 at T6 condition. Figure 2 The image shows a uniformly dispersed precipitated phase; a transmission electron microscope (TEM) image of Example 1 at T6. Figure 3 The image shows uniformly dispersed Ω and θ' phases; a scanning electron microscope (SEM) image of the ZL205A alloy in the T6 condition of Comparative Example 1. Figure 4 It shows coarse, segregated precipitates.

[0113] As can be seen from the above, compared with the comparative examples, the embodiments of the present invention, through the organic combination of multi-element microalloying composition design, low differential pressure casting process, and multi-stage heat treatment technology, successfully prepared cast aluminum alloys with high room temperature strength, high plasticity, excellent high temperature strength, and fine grain structure. This overcomes the shortcomings of existing technologies, such as the difficulty in simultaneously achieving high temperature strength and room temperature plasticity, coarse microstructure, and the harmful effects of impurity phases. It has significant industrial practical value and broad application prospects. Furthermore, it can be seen that the overall effect is even better when all process parameters are within the preferred range of the present invention.

[0114] 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 high-strength, high-toughness, heat-resistant cast aluminum alloy, characterized in that, The composition, by weight percentage, includes Cu 4.5-5.5%, Mn 0.4-0.8%, Mg 0.15-0.3%, Zr 0.15-0.25%, V 0.08-0.15%, Ti 0.08-0.15%, B 0.005-0.02%, Cd 0.15-0.25%, Ag 0.05-0.20%, RE 0.05-0.20%, Fe 0.01-0.05%, Si 0.01-0.03%, with the balance being Al and unavoidable impurity elements, wherein each individual impurity element is ≤0.03% and the total amount of the impurity elements is ≤0.12%.

2. The high-strength, high-toughness, heat-resistant cast aluminum alloy according to claim 1, characterized in that, The RE comprises La and Ce, and the weight ratio of La to Ce is 1:(1~3); and / or The weight ratio of Mn to Cu is 0.08~0.15; and / or The weight ratio of Zr to Ti is 1.2 to 2.0; and / or The weight ratio of Ti to B is 6 to 26.

3. The high-strength, high-toughness, heat-resistant cast aluminum alloy according to claim 1 or 2, characterized in that, By weight percentage, the high-strength, high-toughness, and heat-resistant cast aluminum alloy comprises Cu 4.8~5.2%, Mn 0.5~0.7%, Mg 0.2~0.28%, Zr 0.18~0.24%, V 0.10~0.14%, Ti 0.09~0.14%, B 0.006~0.015%, Cd 0.18~0.24%, Ag 0.08~0.15%, RE 0.08~0.18%, Fe 0.01~0.05%, Si 0.01~0.03%, with the balance being Al and unavoidable impurity elements.

4. The method for preparing the high-strength, high-toughness, and heat-resistant cast aluminum alloy according to any one of claims 1 to 3, characterized in that, Includes the following steps: Step S1: According to the composition of the high-strength, tough, and heat-resistant cast aluminum alloy, the raw materials are batched, smelted, and refined in sequence to obtain the aluminum alloy melt; Step S2: The aluminum alloy melt is transferred to a casting machine for casting, the casting including low-pressure casting or differential pressure casting, to obtain a casting blank. Step S3: The cast blank is subjected to two-stage solution treatment, quenching and two-stage aging in sequence to obtain the high-strength, high-toughness and heat-resistant cast aluminum alloy. The two-stage solution treatment includes a primary solution treatment and a secondary solution treatment, wherein the temperature of the primary solution treatment is less than the temperature of the secondary solution treatment, with a difference of 22~38℃; the two-stage aging treatment includes a primary aging treatment and a secondary aging treatment, wherein the temperature of the primary aging treatment is less than the temperature of the secondary aging treatment, with a difference of 40~65℃.

5. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 4, characterized in that, In step S2 The casting machine is a low-pressure casting machine or a differential pressure casting machine; and / or The casting temperature is 680~700℃; and / or Step S2 includes: placing the aluminum alloy melt in the holding furnace of the casting machine, connecting the aluminum alloy melt to the mold through a riser pipe, introducing dry compressed air to create pressure on the surface of the aluminum alloy melt, and causing the aluminum alloy melt to sequentially rise, fill, pressurize, hold, and depressurize along the riser pipe from bottom to top to obtain the casting blank.

6. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 5, characterized in that, In step S2 The riser pipe diameter of the casting machine is 90~160mm; and / or The low-pressure casting pressure is 110~160 kPa, the liquid rise rate is 1.2~1.7 kPa / s, the mold filling rate is 0.8~1.1 kPa / s, the crust formation pressure increase value is 10~20 kPa, the crust formation time is 10~20 s, and the pressure holding time is 200~800 s; and / or The differential pressure casting process involves a lower tank pressure of 400-500 kPa, an upper tank pressure of 300-400 kPa, a pressure difference of 100-200 kPa, a liquid rise rate of 1.1-1.6 kPa / s, a filling rate of 0.9-1.1 kPa / s, a crust formation pressure increase of 12-18 kPa, a crust formation time of 10-20 s, and a holding time of 300-600 s.

7. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 4, characterized in that, In step S3 The temperature of the first solution treatment is 500~520℃, and the holding time is 4~8h; and / or The secondary solution treatment is performed at a temperature of 535~545℃ for 8~12 hours; and / or The quenching is water quenching at a temperature of 60~80℃ and a quenching transfer time of 5~10s.

8. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 4, characterized in that, In step S3 The temperature for the first aging process is 100~120℃, and the holding time is 4~6 hours; and / or The temperature for the secondary aging process is 150~170℃, and the holding time is 10~14h.

9. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 4, characterized in that, Step S1 includes: Step S11: Prepare materials according to the composition of the high-strength, tough, and heat-resistant cast aluminum alloy to obtain aluminum ingots, AlCu master alloy, AlMn master alloy, Mg ingots, AlZr master alloy, AlV master alloy, AlTi master alloy, AlB master alloy, AlCd master alloy, AlAg master alloy, and AlRE master alloy. Step S12: The aluminum ingot is subjected to a first heat preservation, and the AlCu master alloy, the AlMn master alloy, the AlV master alloy and the AlCd master alloy are added sequentially at the first heat preservation temperature to obtain a first intermediate melt; Step S13: The first intermediate melt is subjected to a second heat preservation. At the second heat preservation temperature, the AlZr master alloy, the AlAg master alloy, the Mg ingot and the AlRE master alloy are added in sequence. After complete melting, the AlTi master alloy and the AlB master alloy are added to obtain the second intermediate melt. Step S14: Argon gas is introduced into the second intermediate melt for refining, and a refining agent is added to obtain the aluminum alloy melt.

10. The method for preparing high-strength, high-toughness, and heat-resistant cast aluminum alloy according to claim 9, characterized in that, The temperature of the first insulation is 720~740℃; and / or The second insulation temperature is 700~720℃; and / or The refining temperature is 710~730℃, the time is 15~25min, and the rotation speed is 500~700r / min.

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