A heat-resistant Al-Cu-Mg-Mn aluminum alloy and a preparation method and application thereof

Abstract

This invention belongs to the field of aluminum alloy hot working and heat treatment technology, specifically disclosing a heat-resistant Al-Cu-Mg-Mn aluminum alloy, its preparation method, and its applications. The aluminum alloy composition of this invention includes Cu, Mg, Mn, Zn, Cr, Zr, and Ti, with the balance being Al, targeting medium-temperature load-bearing applications at 250–300℃. The process of this invention does not rely on adding high-cost elements or rapid solidification and additive manufacturing as the main technical route. Instead, it introduces a pre-forging process after the alloy casting billet is completed and before homogenization treatment. This process breaks down and reconstructs the continuous coarse second-phase skeleton in the casting state, forming dislocations, subgrain boundaries, and local high-strain zones in the matrix. Subsequently, through graded homogenization, hot extrusion, two-stage solution treatment, and artificial aging, high-density, fine, and uniform precipitation of the T-phase is induced. This method can be integrated with conventional casting, forging, extrusion, and heat treatment equipment, and is suitable for medium-temperature heat-resistant load-bearing components in aerospace and transportation equipment.

Description

Technical Field

[0001] This invention relates to the field of aluminum alloy hot working and heat treatment technology, and in particular to a heat-resistant Al-Cu-Mg-Mn aluminum alloy, its preparation method and application. Background Technology

[0002] Aluminum alloys, due to their low density, high specific strength, and good machinability, have significant application value in aerospace, transportation, and high-end equipment. For applications such as supersonic aircraft, hot-end connectors, and lightweight components near power systems, materials must maintain load-bearing capacity and microstructural stability at temperatures of 250–300°C or even higher. Traditional 2xxx series Al-Cu or Al-Cu-Mg wrought aluminum alloys primarily rely on age-induced precipitates such as θ′-Al2Cu and S-Al2CuMg phases for strengthening. However, these precipitates are prone to coarsening, dissolution, or migration to grain boundaries under high temperatures or long-term thermal exposure, leading to weakened intragranular strengthening, continuous grain boundary precipitation, and widening of the non-precipitate zone.

[0003] Currently, research on heat-resistant aluminum alloys mainly follows two routes. The first is microalloying, which involves adding elements such as Sc, Ag, Er, Zr, Y, Ce, and Ni to alloys like Al-Cu, Al-Mg, and Al-Si to improve high-temperature stability by forming Ω phases, L12 phases, interfacial segregation structures, or rare-earth / transition metal-rich phases. However, this route relies on high-cost elements, and some systems suffer from problems such as long heat treatment times, limited precipitate quantities, and difficulty in simultaneously achieving room-temperature strength and high-temperature stability. The second route involves eutectic / near-eutectic heat-resistant aluminum alloys assisted by rapid solidification, powder metallurgy, or additive manufacturing. These alloys form fine Al-Fe, Al-Ni, Al-Ce, or Al-Si-based heat-resistant second phases through high cooling rates, exhibiting better microstructural stability. However, this method is highly dependent on powder quality, solidification rate, forming window, and densification process, and remains limited in terms of size, shape, defect control, and engineering scale-up. Therefore, for aerospace sheet metal, extrusions, and forgings, there is still a need to develop low-cost microstructure control methods that are compatible with traditional casting-hot working-heat treatment processes.

[0004] In Al-Cu-Mn or Al-Cu-Mg-Mn systems, the T-Al20Cu2Mn3 phase is a Cu- and Mn-rich heat-resistant phase with high thermal stability, and is expected to provide continuous resistance to dislocation movement and grain boundary stability under medium-temperature service conditions of 250–300℃. Existing technologies focus more on controlling the T phase through composition design, trace element segregation, or aging regimes. However, in conventional casting and hot working processes, the as-cast coarse second phase often presents as a continuous skeleton or network distribution with significant local segregation. This results in limited nucleation sites for the T phase during subsequent homogenization and aging processes, easily leading to insufficient numbers and excessively large T phase particles.

[0005] Therefore, there is an urgent need for a preparation method that does not rely on expensive elements such as Sc / Ag, does not rely on special equipment for rapid solidification or additive manufacturing, and can significantly improve the precipitation density and uniformity of T phase in Al-Cu-Mg-Mn alloys, so as to improve the high-temperature strength of this type of aluminum alloy under service conditions of 250-300℃ and the ability to retain strength after heat exposure. Summary of the Invention

[0006] This invention aims to at least solve one of the technical problems existing in the prior art. To this end, this invention provides a method for preparing heat-resistant Al-Cu-Mg-Mn aluminum alloys based on pre-forging before homogenization to induce high-density precipitation of the T-Al2OCu2Mn3 phase, as well as the aluminum alloy material obtained by this method and its applications. This solves the problems of coarsening of precipitated phases, continuous grain boundary precipitation, insufficient T-phase number density, and insufficient strength retention after heat exposure in existing Al-Cu-Mg or Al-Cu-Mg-Mn heat-resistant aluminum alloys under service conditions of 250–300℃.

[0007] Traditional methods involve homogenizing cast aluminum alloy billets before hot working to reduce segregation, dissolve low-melting-point phases, and improve hot working plasticity. This invention, through exploration, reveals an unconventional process: pre-forging the cast billet before homogenization. Before the as-cast microstructure is fully eliminated by traditional homogenization, controlled plastic deformation breaks down continuous coarse phases, shortens solute diffusion distance, and establishes numerous heterogeneous nucleation sites. This alters the subsequent T-phase precipitation path, achieving a transformation from a small amount of coarse precipitation to high-density refined precipitation of the T-phase.

[0008] The core concept of this invention lies in the fact that, instead of simply increasing the content of expensive microalloying elements or using rapid solidification / additive manufacturing to obtain a refined microstructure, a controlled pre-forging step is introduced after conventional casting and before homogenization. Pre-forging first mechanically breaks down and spatially reconstructs the as-cast coarse second-phase continuous skeleton, simultaneously forming dislocations, subgrain boundaries, local orientation differences, and high-strain energy storage regions within the matrix. Subsequently, graded homogenization promotes solute redistribution and the spheroidization / dissolution / redispersal of the broken coarse phase, resulting in more heterogeneous nucleation sites and shorter diffusion distances for the T-Al₂OCu₂Mn₃ phase during subsequent hot extrusion, solution treatment, and aging processes, ultimately forming a larger number of smaller, more uniformly distributed T-phases.

[0009] In a first aspect, the present invention provides a heat-resistant Al-Cu-Mg-Mn aluminum alloy, wherein, by mass percentage, the alloy composition comprises Cu 4.3-4.6%, Mg 1.5-1.7%, Mn 0.4-0.5%, Zn 0.5-0.6%, Cr 0.20-0.25%, Zr 0.15-0.20%, Ti 0.08-0.10%, with the balance being Al and unavoidable impurities.

[0010] According to specific embodiments of the present invention, the heat-resistant aluminum alloy provided by the present invention is designed for medium-temperature load-bearing applications at 250–300°C, without the need to add high-cost elements such as Sc, Ag, and Er. In the Al-Cu-Mg-Mn aluminum alloy of the present invention, the T-Al2OCu2Mn3 phase is highly dense, fine, and uniformly precipitated, with a small average T phase size and a higher number of particles per unit area. This aluminum alloy exhibits significantly improved tensile strength at 300°C and room temperature tensile strength after 100 hours of heat exposure at 300°C, making it particularly suitable for medium-temperature heat-resistant load-bearing components in aerospace and transportation equipment.

[0011] In the Al-Cu-Mg-Mn aluminum alloy of the present invention, Cu and Mg provide the basis for strengthening the Al-Cu-Mg matrix, Mn provides the necessary element for the formation of the T-Al20Cu2Mn3 phase, Cr, Zr and Ti are used to assist in the stability of grain and hot working structure, and Zn is used to adjust the supersaturated solid solution state and the aging precipitation response.

[0012] According to some embodiments of the present invention, the aluminum alloy contains a dispersed T-Al20Cu2Mn3 phase in the T6 state, the average size of the T-Al20Cu2Mn3 phase being 120–180 nm, and the maximum size within 1 μm. 2 The average number of T-Al20Cu2Mn3 phase particles within the statistical area is 5 to 8.

[0013] According to some embodiments of the present invention, at 135.2 μm 2 The average number of T-Al20Cu2Mn3 phase particles within the statistical area is 700–950, which translates to approximately 1 μm. 2 The average number of particles within the statistical area was 5.2 to 7.0.

[0014] According to some embodiments of the present invention, the aluminum alloy is an aluminum alloy material that has undergone hot deformation processing and T6 heat treatment, and its tensile strength in a high-temperature tensile test at 300℃ according to GB / T 228.2-2015 is not less than 260 MPa; and / or, after being exposed to heat at 300℃ for 100 h, its tensile strength in a room-temperature tensile test according to GB / T 228.1-2021 is not less than 270 MPa; wherein, the tensile specimen is a round bar specimen cut along the hot deformation direction, with a gauge length diameter of 3 mm and a gauge length of 15 mm.

[0015] A second aspect of the present invention provides a method for preparing the heat-resistant Al-Cu-Mg-Mn aluminum alloy described in the first aspect of the present invention, comprising the following steps:

[0016] S1. Al, Mg, Zn and Al-Cu, Al-Mn, Al-Cr, Al-Zr and Al-Ti intermediate alloy raw materials are prepared according to the mass percentage of alloy components, mixed and smelted, then refined, stirred and cast to obtain cast billets.

[0017] S2. Heat the cast billet to 430-470℃ for pre-forging treatment;

[0018] S3. The pre-forged billet is graded and homogenized, and then subjected to hot deformation processing, solution treatment, quenching and aging treatment to obtain a heat-resistant Al-Cu-Mg-Mn aluminum alloy containing high-density T-Al20Cu2Mn3 phase.

[0019] The method for preparing heat-resistant Al-Cu-Mg-Mn aluminum alloys of the present invention does not rely on adding high-cost elements such as Sc, Ag, and Er, or on rapid solidification and additive manufacturing as the main technical route. Instead, it introduces a pre-forging process after the Al-Cu-Mg-Mn alloy as-cast billet is completed and before the homogenization process, so that the continuous coarse second-phase skeleton in the as-cast state is broken and reconstructed, and dislocations, subgrain boundaries and local high-strain regions are formed in the matrix. Subsequently, through graded homogenization, hot extrusion, two-stage solid solution and artificial aging, high-density, fine and uniform precipitation of T-Al20Cu2Mn3 phase is induced.

[0020] According to some embodiments of the present invention, in step S1, the melting temperature is 720-740°C.

[0021] According to some embodiments of the present invention, in step S1, the refining is performed once or twice.

[0022] According to some embodiments of the present invention, in step S2, the pre-forging process is a combination of multi-pass upsetting and drawing deformation.

[0023] According to some embodiments of the present invention, the pre-forging is a four-upsetting and three-drawing forging system, which includes adjusting the stress direction of the billet between each upsetting or drawing pass, so that the billet is deformed alternately in the height and length directions; the pre-forging temperature is 440 to 460°C, and the single-pass reduction is 15% to 25%.

[0024] The four-upsetting and three-drawing process of this invention, through alternating upsetting and drawing, subjectes the continuous second phase between dendrites in the as-cast state to combined shearing, compression, and stretching in multiple directions. This is more conducive to breaking the continuous skeleton and shortening the solute diffusion distance than unidirectional compression or single-pass deformation, and to forming more uniform dislocations, subgrain boundaries, and local high-strain regions in the matrix. In this step, the coarse Al-Cu-Mg phase, Al-Cu-Mn phase, and Fe / Mn-containing second phase formed by unavoidable impurities that are continuously distributed along the dendrites or near the grain boundaries in the as-cast state are cut off, crushed, and redispersed. At the same time, high-density dislocations, subgrain boundaries, and local high-strain energy storage regions are formed inside the aluminum matrix. Fe is not an actively added element, but only comes from unavoidable impurities in the aluminum alloy raw materials or intermediate alloys.

[0025] Among these parameters, pre-forging temperature and single-pass reduction are crucial for influencing the subsequent precipitation density of the T-phase. If the pre-forging temperature is too low, the as-cast billet lacks sufficient thermoplasticity, making it difficult to fully break down the coarse second phase and prone to cracking. If the pre-forging temperature is too high, the as-cast low-melting-point segregation region carries the risk of localized overheating or grain boundary weakening, and may also weaken the deformation energy storage and defect retention effects. When the single-pass reduction is below 15%, the breakage of the coarse continuous phase and the introduction of matrix defects are insufficient; when the single-pass reduction is above 25%, the risk of billet cracking and localized microstructure inhomogeneity increases. Therefore, a single-pass reduction of 440–460℃ and 15%–25% is beneficial for achieving a balance between hot working safety and subsequent high-density T-phase precipitation.

[0026] The pre-forging process specially designed in this invention is set between the cast billet and the graded homogenization process, and the cast billet is not subjected to conventional high-temperature long-term homogenization process before the pre-forging process.

[0027] According to some embodiments of the present invention, in step S3, the first stage of the graded homogenization treatment is to keep the temperature at 460-480℃ for 20-28 hours, and the second stage is to keep the temperature at 480-500℃ for 20-28 hours.

[0028] This invention utilizes a first-stage, lower-temperature homogenization process to reduce the risk of overheating or localized melting while promoting the spheroidization, discontinuity, and initial dissolution of the as-cast low-melting-point eutectic structure and the fragmented second phase. The second-stage, higher-temperature homogenization further enhances the diffusion and solid solution of solutes such as Cu, Mg, and Mn, allowing the phase interfaces, dislocations, and subgrain boundaries formed during pre-forging to transform into effective heterogeneous nucleation sites for the T-phase during subsequent solid solution and aging processes. Compared to single high-temperature direct homogenization, the two-stage homogenization method employed in this invention is more advantageous in balancing hot working safety, solute homogenization, and high-density T-phase precipitation.

[0029] According to some embodiments of the present invention, in step S3, the hot deformation process is a hot extrusion process, the hot extrusion temperature is 430~470℃, and the extrusion ratio is (20~30):1.

[0030] According to some embodiments of the present invention, in step S3, the solution treatment is a two-stage solution treatment, first at 495-505℃ for 0.5-1.5 h, and then at 505-515℃ for 0.5-1.5 h.

[0031] According to some embodiments of the present invention, in step S3, the aging process involves maintaining a temperature of 150–170°C for 16–24 hours.

[0032] The above-mentioned processes of the present invention are not isolated from each other, but form a synergistic relationship with pre-forging before homogenization: pre-forging establishes a basis for controllable defects and phase reconstruction, graded homogenization realizes solute redistribution, hot extrusion further refines and densifies the structure, and two-stage solid solution and aging ultimately induce the precipitation of high-density T phase.

[0033] In a third aspect, the present invention provides the application of the heat-resistant Al-Cu-Mg-Mn aluminum alloy described in the first aspect of the present invention in medium-temperature load-bearing components under service conditions of 250 to 300°C.

[0034] According to some embodiments of the present invention, the intermediate-temperature load-bearing component includes at least one of aerospace structural components, hot-end adjacent connectors, heat-resistant fastening load-bearing components, missile or aircraft shell / wing surface components, and lightweight heat-resistant components for transportation equipment.

[0035] The beneficial effects of this invention are:

[0036] 1) This invention sets the pre-forging process before homogenization, and actively reconstructs the coarse second phase in the as-cast state using an unconventional process sequence, rather than passively waiting for homogenization to eliminate segregation, thereby increasing the number of nucleation sites for subsequent precipitation of the T phase from the source.

[0037] 2) This invention achieves heat resistance strengthening through the T-Al20Cu2Mn3 phase in the low-cost Al-Cu-Mg-Mn alloy system, without relying on high-cost elements such as Sc, Ag, and Er as the main strengthening agents, thus reducing material costs and the difficulty of engineering promotion.

[0038] 3) This invention can be integrated with conventional casting, forging, extrusion and heat treatment equipment, and does not require fast solidification powder, special additive manufacturing powder and complex forming windows, making it suitable for scaling up to larger forgings or extrusions.

[0039] 4) The number of T phases obtained by the present invention is nearly twice that of the unforged comparative example, and the size is significantly refined, which can improve the instantaneous tensile strength at 300℃ and the room temperature strength retention ability after 300℃ heat exposure.

[0040] Other features and advantages of the invention will be set forth in the description which follows, and will be apparent in part from the description, or may be learned by practicing the invention. Attached Figure Description

[0041] The present invention will be further described below with reference to the accompanying drawings and embodiments, wherein:

[0042] Figure 1 This is a schematic diagram of the preparation process of the heat-resistant aluminum alloy in Embodiment 1 of the present invention;

[0043] Figure 2 This is a comparative diagram showing the morphology and quantity of the T phase in the T6 state of the aluminum alloy samples prepared in Example 1 and Comparative Example 1 of this invention.

[0044] Figure 3 The figures show a comparison of the mechanical properties of aluminum alloy samples prepared in Example 1 and Comparative Example 1 of the present invention. Figure a shows the stress-strain curves of room temperature tensile test in the T6 state and room temperature tensile test after 100 h of heat exposure at 300℃. Figure b shows the stress-strain curve of high temperature tensile test at 300℃.

[0045] Figure 4 The image shows a comparison of the grain boundary microstructure of aluminum alloy samples prepared in Example 1 and Comparative Example 1 after being exposed to heat at 300°C for 100 h. Detailed Implementation

[0046] The following will describe the concept and technical effects of the present invention clearly and completely with reference to embodiments, so as to fully understand the purpose, features and effects of the present invention. Obviously, the described embodiments are only some embodiments of the present invention, not all embodiments. Other embodiments obtained by those skilled in the art based on the embodiments of the present invention without creative effort are all within the scope of protection of the present invention.

[0047] Unless otherwise specified in the examples, the procedures should be performed under standard conditions or conditions recommended by the manufacturer. Reagents or instruments whose manufacturers are not specified are all commercially available products.

[0048] Example 1

[0049] This embodiment provides the preparation of Al-Cu-Mg-Mn heat-resistant aluminum alloys.

[0050] This embodiment is based on a process flow of pre-forging before homogenization—graded homogenization—hot extrusion—two-stage solution treatment—artificial aging, as shown in the schematic diagram below. Figure 1 As shown, the specific steps are as follows:

[0051] 1) Prepare Al-Cu-Mg-Mn heat-resistant aluminum alloys according to the composition shown in Table 1. Add pure Al, pure Mg, pure Zn and Al-50Cu, Al-20Mn, Al-10Cr, Al-10Zr and Al-10Ti master alloys to the melting equipment, melt at about 730℃ and hold at that temperature.

[0052] 2) After the alloying elements have been fully melted, the mixture is refined twice and then stirred with electromagnetic stirring to improve the uniformity of the melt composition. The mixture is then poured into a preheated copper mold to obtain a cast billet.

[0053] 3) The obtained cast billet is directly heated to 450℃ and held for 60 min before homogenization pre-forging. The pre-forging adopts a four-upsetting and three-drawing system, that is, multiple-pass combined deformation is carried out in the order of "upsetting - drawing - upsetting - drawing - upsetting - drawing - upsetting - drawing - upsetting". After each deformation is completed, the stress direction of the billet is adjusted, and the billet is returned to the furnace for holding for 10-15 min according to the temperature drop of the billet. The single-pass reduction is 20%, and the final forging temperature is controlled above 400℃.

[0054] The above-mentioned four-upsetting and three-drawing system, through alternating upsetting and drawing, causes the continuous second phase between the as-cast dendrites to be subjected to combined shearing, compression and stretching in multiple directions. This is more conducive to breaking the continuous skeleton and shortening the solute diffusion distance than unidirectional compression or single-pass deformation, and to forming more uniform dislocations, subgrain boundaries and local high-strain regions in the matrix.

[0055] In this step, the coarse Al-Cu-Mg phase, Al-Cu-Mn phase, and Fe / Mn-containing second phase formed by unavoidable impurities that are continuously distributed along the dendrites or near the grain boundaries in the as-cast state are cut off, crushed, and redispersed; at the same time, high-density dislocations, subgrain boundaries, and local high-strain energy storage regions are formed inside the aluminum matrix; Fe is not an active additive element, but only comes from unavoidable impurities in the aluminum alloy raw materials or intermediate alloys;

[0056] 4) After pre-forging, a two-stage homogenization treatment is carried out, with the homogenization treatment regime being 470℃ for 24 h and 490℃ for 24 h;

[0057] Compared with direct homogenization, the two-stage homogenization process in this embodiment can reduce the size and continuity of the coarse second phase after pre-forging, shorten the solute diffusion distance during the homogenization process, and make the local redistribution of elements such as Cu, Mn, and Mg more sufficient, which is conducive to the simultaneous nucleation of the T phase in more locations.

[0058] 5) The homogenized billet was hot-extruded at 450℃ with an extrusion ratio of 25:1; after extrusion, it was subjected to two-stage solution treatment at 500℃ for 1 h and 510℃ for 1 h, followed by water quenching; and then artificial aging at 160℃ for 20 h to obtain the heat-resistant aluminum alloy, which was designated as PHF sample.

[0059] Hot extrusion further breaks down the residual coarse second phase, improves the compactness of the billet, and forms a more uniform deformed structure.

[0060] Comparative Example 1

[0061] This comparative example provides the preparation of an Al-Cu-Mg-Mn aluminum alloy. This comparative example is essentially the same as Example 1, employing the same chemical composition, melting and casting, two-stage homogenization, hot extrusion, two-stage solution treatment, and artificial aging process. However, the pre-forging step before homogenization is omitted to obtain the aluminum alloy, denoted as the Base sample.

[0062] Comparative Example 2

[0063] This comparative example provides the preparation of Al-Cu-Mg-Mn aluminum alloy. This comparative example is essentially the same as Example 1, except that in step 3), the pre-forging temperature is 400℃. All other steps and processes are consistent with Example 1. The final aluminum alloy sample is obtained.

[0064] Comparative Example 3

[0065] This comparative example provides the preparation of an Al-Cu-Mg-Mn aluminum alloy. This comparative example is essentially the same as Example 1, except that in step 3), the pre-forging temperature is 500°C; all other steps and processes are consistent with Example 1. The final aluminum alloy sample is obtained.

[0066] Comparative Example 4

[0067] This comparative example provides the preparation of Al-Cu-Mg-Mn aluminum alloy. This comparative example is basically the same as Example 1, except that in step 3), the single-pass reduction is 10%, while other steps and processes are consistent with Example 1. The final aluminum alloy sample is obtained.

[0068] Comparative Example 5

[0069] This comparative example provides the preparation of Al-Cu-Mg-Mn aluminum alloy. This comparative example is basically the same as Example 1, except that in step 3), the single-pass reduction is 30%, while other steps and processes are consistent with Example 1. The final aluminum alloy sample is obtained.

[0070] Organizational Analysis:

[0071] Tissue analysis was performed on the PHF sample from Example 1 and the Comparative Example 1 (Base sample), and the results are as follows: Figure 2 As shown.

[0072] Microstructural analysis showed that the basic phase composition of the PHF sample in Example 1 and the Comparative Example 1 (Base sample) both included an α-Al matrix and phases such as Al2Cu, Al2CuMg, T-Al20Cu2Mn3, and Al7Cu2Mn, or Al7Cu2(Fe, Mn) formed by unavoidable Fe impurities, indicating that pre-forging before homogenization did not introduce new undesirable phases. However, there were significant differences in phase size, phase continuity, and T-phase precipitation density between the two samples.

[0073] In the T6 state, the T phase in Comparative Example 1 (Base sample) is more likely to be blocky or short and thick rod-shaped, and its quantity is relatively limited; in the PHF sample of Example 1, the T phase is mainly distributed in small short rod-shaped or nearly spherical shapes. Statistical results show that the average size of the T phase decreased from about 233 nm in Comparative Example 1 to about 148 nm in Example 1; at the same 135.2 μm... 2 Within the statistical area, the number of T-phase particles increased from approximately 444 in Comparative Example 1 to approximately 851 in Example 1, which is approximately 1.9 times that of the Comparative Example.

[0074] The above results demonstrate that pre-forging before homogenization can transform the coarse, continuous network of the as-cast second phase into more discrete granular or short rod-shaped residual phases, and provide heterogeneous nucleation sites in the matrix. This changes the T-phase precipitation pattern from "few nucleations, continuous growth" to "high-density nucleation, restricted growth," thereby simultaneously increasing the number and refining the size of the T-phase. Comparative Example 1 eliminates differences in composition and subsequent heat treatment, allowing direct evaluation of the contribution of pre-forging before homogenization to T-phase precipitation and high-temperature performance.

[0075] In Comparative Example 2, due to insufficient thermoplasticity of the cast billet, the coarse second phase was not sufficiently broken down, and the risk of local cracking increased during deformation. The final product had an average T-phase size of approximately 205 nm, with a diameter of 135.2 μm. 2 The number of particles within the statistical area was approximately 560. In Comparative Example 3, at this temperature, the grain boundary segregation region was prone to local overheating or coarsening, and the deformation energy storage was insufficient. The average size of the T phase in the final product was approximately 190 nm, with a diameter of 135.2 μm. 2 The number of particles within the statistical area was approximately 615; in Comparative Example 4, due to insufficient total deformation and fragmentation, a larger amount of the continuous second phase was retained, with an average T-phase size of approximately 214 nm and a diameter of 135.2 μm. 2 The number of particles within the statistical area was approximately 535. In Comparative Example 5, although the degree of fragmentation of the coarse second phase was improved, the local strain concentration and tendency for microcracks were enhanced, and the uniformity of the microstructure decreased. The average size of the T phase was approximately 172 nm, with a diameter of 135.2 μm. 2 The number of particles within the statistical area is approximately 690.

[0076] Further analysis of Comparative Examples 2-5 reveals that pre-forging does not achieve the same effect at any temperature or with any amount of deformation. When the pre-forging temperature is too low or the single-pass reduction is insufficient, the continuous second phase in the as-cast state is difficult to fully break down, resulting in a limited number of T-phase nucleation sites. When the pre-forging temperature is too high or the single-pass reduction is too large, although it can promote the discontinuity of the second phase, it easily leads to local overheating, recovery softening, strain concentration, or a tendency for microcracks, weakening the subsequent T-phase refinement and performance retention effects. Therefore, the pre-forging regime before homogenization defined in this invention exhibits a clear process window characteristic.

[0077] The above results indicate that the process of the present invention can achieve a better balance between hot working safety, multi-directional crushing effect of the second phase, deformation energy storage retention and high-density precipitation of the T phase.

[0078] Performance testing:

[0079] Regarding mechanical properties, the comparison results between Example 1 and Comparative Example 1 are shown in Table 2. Room temperature tensile testing was conducted according to GB / T228.1-2021, and high-temperature tensile testing at 300℃ was conducted according to GB / T 228.2-2015. The heat exposure treatment involved holding the T6 state samples at 300℃ for 100 h, cooling them to room temperature, and then conducting room temperature tensile testing. Tensile specimens were cut along the extrusion direction, using round bar tensile specimens with a gauge length diameter of 3 mm and a gauge length of 15 mm, for a total length of approximately 45 mm. Each group contained 3 valid specimens, and the average test result was taken.

[0080] From Table 2 and Figure 3 visible, Figure 3 Figure a shows the room temperature tensile curves of the aluminum alloy samples in the T6 state and the room temperature tensile curves after heat exposure at 300°C for 100 h for Example 1 and Comparative Example 1. Figure 3 Figure b shows the high-temperature tensile curves at 300℃ for Example 1 and Comparative Example 1. The room-temperature tensile strength of Example 1 in the T6 state was 521.9 MPa, slightly lower than the 544.5 MPa of Comparative Example 1, but the elongation in the T6 state increased from 16.49% in Comparative Example 1 to 18.17%, indicating that pre-forging before homogenization did not sacrifice the room-temperature plasticity of the material. More importantly, under the high-temperature tensile conditions at 300℃, the tensile strength of Example 1 reached 272.9 MPa, significantly higher than the 229.4 MPa of Comparative Example 1, an increase of approximately 19.0%; after 100 h of heat exposure at 300℃, the room-temperature tensile strength of Example 1 remained at 277.2 MPa, higher than the 237.3 MPa of Comparative Example 1, an increase of approximately 16.8%. Therefore, from... Figure 3The results show that the performance of the present invention varies in three dimensions: T6 state, 300℃ high temperature tensile test, and residual strength after 300℃ heat exposure. The most prominent technical effect is reflected in the instantaneous load-bearing capacity at medium temperature and the strength retention capacity after heat exposure.

[0081] Therefore, the role of pre-forging before homogenization is not simply to increase the peak room temperature strength of the T6 state, but rather to significantly improve the material's load-bearing capacity at 300℃ and its strength retention after long-term heat exposure at 300℃, while maintaining high room temperature strength and good elongation. Combined with... Figure 3 The stress-strain curves show that Example 1 exhibits higher rheological stress and tensile strength during the high-temperature tensile stage, and still possesses higher residual room-temperature strength after heat exposure, indicating that this process is more in line with the service requirements of medium-temperature heat-resistant components at 250–300℃. This performance change is related to the increased number, finer size, and more uniform distribution of the T-Al20Cu2Mn3 phase, the reduced continuity of grain boundary precipitation, and the suppression of PFZ width after heat exposure. Compared with Comparative Example 1 (Base sample), the strength advantage of Example 1 under high-temperature and heat exposure conditions demonstrates that the high-density precipitation of the T phase induced by pre-forging can continue to provide relatively stable dislocation pinning and microstructure stabilization near grain boundaries after the θ′ phase and S phase coarsen or dissolve back.

[0082] Combination Figure 4 It is evident that after 100 h of heat exposure at 300℃, the precipitates near the grain boundaries in Comparative Example 1 (Base sample) are more prone to coarsening and continuous distribution, with more pronounced no-precipitate bands on both sides of the grain boundaries, indicating that solute depletion and precipitate migration / coarsening are more severe near the grain boundaries during heat exposure. The continuity of grain boundary precipitates in Example 1 sample is weaker than that in the Comparative Example, with the width of the no-precipitate band near the grain boundaries preferably being 100–160 nm, and a large number of fine, dispersed T-phase particles still remaining within the grains. These differences in microstructure indicate that pre-forging before homogenization breaks down the continuous coarse phases in the as-cast state and increases the nucleation sites within the grains, allowing more Cu and Mn-related heat-resistant phases to disperse and precipitate within the grains, reducing the tendency for continuous precipitation at the grain boundaries. Simultaneously, the T-Al20Cu2Mn3 phase has higher thermal stability than the θ′ phase and S phase, and can continuously hinder dislocation movement and local softening near the grain boundaries during heat exposure. Therefore, Example 1 exhibits higher residual strength after heat exposure.

[0083] From a process scale-up perspective, this invention does not require powder metallurgy, rapid solidification, or additive manufacturing processes, nor does it require the addition of large amounts of high-cost elements such as Sc and Ag. Instead, it adds a controllable pre-forging step before homogenization to conventional cast billets. Therefore, this invention is suitable for integration with existing aluminum alloy forging, extrusion, and heat treatment production lines, and has good engineering application prospects for medium-temperature load-bearing aluminum alloy components.

[0084] The embodiments of the present invention have been described in detail above. However, the present invention is not limited to the above embodiments. Within the scope of knowledge possessed by those skilled in the art, various changes can be made without departing from the spirit of the present invention. Furthermore, the embodiments of the present invention and the features thereof can be combined with each other unless otherwise specified.

Claims

1. A heat-resistant Al-Cu-Mg-Mn aluminum alloy, characterized in that, By mass percentage, the alloy composition includes Cu 4.3–4.6%, Mg 1.5–1.7%, Mn 0.4–0.5%, Zn 0.5–0.6%, Cr 0.20–0.25%, Zr 0.15–0.20%, Ti 0.08–0.10%, with the balance being Al and unavoidable impurities.

2. The heat-resistant Al-Cu-Mg-Mn aluminum alloy according to claim 1, characterized in that, The aluminum alloy contains a dispersed T-Al₂₀Cu₂Mn₃ phase in the T₆ state, the average size of which is 120–180 nm, and the 1 μm... 2 The average number of T-Al20Cu2Mn3 phase particles within the statistical area is 5 to 8.

3. The heat-resistant Al-Cu-Mg-Mn aluminum alloy according to claim 1, characterized in that, The aluminum alloy is an aluminum alloy material that has undergone hot deformation processing and T6 heat treatment, and its tensile strength in a high-temperature tensile test at 300℃ according to GB / T 228.2-2015 is not less than 260 MPa; and / or, after being exposed to heat at 300℃ for 100 h, its tensile strength in a room-temperature tensile test according to GB / T 228.1-2021 is not less than 270 MPa; wherein, the tensile specimen is a round bar specimen cut along the hot deformation direction, with a gauge length diameter of 3 mm and a gauge length of 15 mm.

4. The method for preparing the heat-resistant Al-Cu-Mg-Mn aluminum alloy according to any one of claims 1 to 3, characterized in that, Includes the following steps: S1. Al, Mg, Zn and Al-Cu, Al-Mn, Al-Cr, Al-Zr and Al-Ti intermediate alloy raw materials are prepared according to the mass percentage of alloy components, mixed and smelted, then refined, stirred and cast to obtain cast billets. S2. Heat the cast billet to 430-470℃ for pre-forging treatment; S3. The pre-forged billet is graded and homogenized, and then subjected to hot deformation processing, solution treatment, quenching and aging treatment to obtain a heat-resistant Al-Cu-Mg-Mn aluminum alloy containing high-density T-Al20Cu2Mn3 phase.

5. The preparation method according to claim 4, characterized in that, In step S1, the melting temperature is 720-740°C; the refining is performed once or twice.

6. The preparation method according to claim 4, characterized in that, In step S2, the pre-forging process is a combination of multi-pass upsetting and drawing deformation.

7. The preparation method according to claim 6, characterized in that, In step S2, the pre-forging is a four-upsetting and three-drawing forging system. The four-upsetting and three-drawing forging system includes adjusting the stress direction of the billet between each upsetting or drawing pass, so that the billet is deformed alternately in the height and length directions. The pre-forging temperature is 440 to 460°C, and the single-pass reduction is 15% to 25%.

8. The preparation method according to claim 4, characterized in that, In step S3, the first stage of the graded homogenization treatment is held at 460-480℃ for 20-28 h, and the second stage is held at 480-500℃ for 20-28 h; the hot deformation processing is hot extrusion processing, the hot extrusion temperature is 430-470℃, and the extrusion ratio is (20-30):

1.

9. The preparation method according to claim 4, characterized in that, In step S3, the solution treatment is a two-stage solution treatment, first at 495-505℃ for 0.5-1.5 h, and then at 505-515℃ for 0.5-1.5 h; the aging is at 150-170℃ for 16-24 h.

10. The application of the heat-resistant Al-Cu-Mg-Mn aluminum alloy as described in any one of claims 1 to 3 in medium-temperature load-bearing components under service conditions of 250 to 300°C; wherein the medium-temperature load-bearing component includes at least one of aerospace structural components, hot-end adjacent connectors, heat-resistant fastening load-bearing components, missile or aircraft shell / wing surface components, and lightweight heat-resistant components for transportation equipment.

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