Al-Cu-Mg-Ag-Mn Series Heat-Resistant Alloy and Its Preparation Method

Abstract

An Al-Cu-Mg-Ag-Mn series heat-resistant alloy and its preparation method. The alloy contains elements and their mass percentages as follows: Cu: 4.8 - 6.4%, Mg: 0.4 - 0.8%, Ag: 0.2 - 1.2%, Mn: 0.2 - 0.8%, Ti: 0 - 0.2%, Mo: 0 - 0.2%, Cr: 0 - 0.2%; the balance is Al, and the sum of the mass percentages of Mo and Cr is 0.2%. The preparation method is: melting and casting, subjecting the ingot to double-stage homogenization annealing at 400 ± 20 °C × (8 - 16) h + 515 ± 15 °C × (12 - 36) h; then hot-rolling a slab with a thickness of 20 - 25 mm to 5 mm at 450 - 480 °C; after annealing at 450 - 480 °C for 1 - 4 h, then cold-rolling to 2 mm; finally, solution treatment at 500 - 550 °C for 1 - 6 h, and then aging at 150 - 180 °C for 8 - 16 h. The obtained alloy has high tensile strength and yield strength at room temperature tensile and 250 °C high-temperature tensile; and when the specimen is thermally exposed at 200 - 250 °C for 100 - 500 h, the alloy still maintains high strength and has good thermal stability.

Description

Technical Field

[0001] The present invention relates to a heat-resistant alloy of Al-Cu-Mg-Ag-Mn system and a preparation method thereof, belonging to the technical field of aluminum alloys. Background Art

[0002] With the rapid development of the national defense industry and the aviation industry, higher requirements are put forward for heat-resistant aluminum alloys. For example, when the speed of a supersonic aircraft increases to 2-4 Mach, the skin temperature of the aircraft will reach 150-250 °C. The aluminum alloy skin is prone to softening under such extreme conditions for a long time, and the strength is significantly reduced. Therefore, developing a new type of heat-resistant aluminum alloy that can be used for a long time at 200-250 °C is an urgent problem to be solved at present.

[0003] It has been found that Ag can change the aging precipitation sequence in the Al-Cu-Mg alloy with high Cu / Mg, and precipitate uniform and fine Ω heat-resistant phases. Ω has better coarsening resistance than the θ′ phase and can exist for a long time without aggregation and growth below 200 °C. In addition, although the addition of Ag significantly improves the high-temperature strength of the Al-Cu-Mg alloy, in addition to being affected by the coarsening of the strengthening phase, matrix softening caused by grain boundary slip and grain coarsening is also an important factor for the strength reduction. Therefore, in order to further improve the thermal strength of the Al-Cu-Mg-Ag alloy, mainly two methods are adopted at present:

[0004] One is fine grain strengthening: that is, to improve the strength of the alloy by refining the grains. However, it is difficult to obtain grains with a size less than 10 μm by traditional mechanical processing methods, and the performance improvement is not obvious. In recent years, obtaining sufficiently fine grains by severe plastic deformation to improve the alloy strength has received wide attention. CN101876041A discloses a preparation method of an Al-Cu-Mg-Ag series ultrafine-grained heat-resistant aluminum alloy, which improves the heat-resistant performance of the Al-Cu-Mg-Ag series alloy in the way of ultrafine grain strengthening by subjecting a homogenized ingot to hot extrusion - solution quenching - equal channel angular pressing - aging treatment. Using this method significantly improves the high-temperature strength of the alloy, but its plasticity is poor, and grain coarsening inevitably occurs at high temperatures. The tissue stability of the nanocrystalline alloy is poor. Once grain growth occurs, it will turn into a general coarse grain structure and lose its excellent performance.

[0005] Second, it is microalloying. By adding rare earth elements such as Er and Ce, the alloy grains are refined. At the same time, trace transition elements such as Mn and Zr are introduced, and a suitable homogenization system is adopted to promote the formation of fine and dispersed phases coherent or semi-coherent with the matrix, which can achieve the effect of dispersion strengthening; they can also inhibit recrystallization, and the pinning of grain boundaries can effectively hinder the slip of grain boundaries at high temperatures and stabilize the grain size. CN105525234A discloses a method for testing the influence of solution aging treatment on the properties of Al-Cu-Mg-Ag alloys containing Zr and Mn, but its research does not involve the formation of dispersed phases and their strengthening mechanisms; CN110724865A discloses a preparation method of an Al-Cu-Mg-Ag-Si-Sc heat-resistant aluminum alloy. By exerting the synergistic effect between Sc and Si elements, Si and Sc elements can segregate around the precipitation phase at a lower aging temperature, inhibiting the coarsening of the precipitation phase, making the alloy have higher strength and heat resistance. However, Sc and Si do not enhance the thermal stability of low-Cu Al-Cu-Mg-Ag alloys. Due to the low diffusion rates of elements such as Mn, Sc, and Zr, during the heat treatment process, due to the different equilibrium distribution coefficients K0, the formation of solute-free zones and sparse zones in the grains occurs, and the distribution is extremely uneven. Therefore, obtaining uniform, fine, and high-number-density dispersed phases is one of the effective ways to improve the high-temperature thermal stability of alloys. However, there are few related studies on how to obtain uniform, fine, and high-number-density dispersed phases in the Al-Cu-Mg-Ag-Mn system. Summary of the Invention

[0006] Aiming at the problems existing in the prior art, the purpose of the present invention is to provide a heat-resistant alloy of the Al-Cu-Mg-Ag-Mn system and its preparation method. Through microalloying with Mo and Cr, and improving the homogenization system, deformation processing technology, and solution and aging heat treatment systems, the synergistic effect between microalloying elements is exerted, promoting the precipitation of spherical Al7(Cr,Mn) and Al6(Fe,Mn,Mo) dispersed phases with a number density of 10.3 / μm 2 , a diameter of 86.2 ± 5 nm and a uniform distribution, so as to improve the dispersion strengthening effect and high-temperature thermal stability of the Al-Cu-Mg-Ag-Mn system heat-resistant alloy.

[0007] A heat-resistant alloy of the Al-Cu-Mg-Ag-Mn system of the present invention includes an Al-Cu-Mg-Ag-Mn matrix and trace amounts of Mo and Cr, and the relationship between Mo and Cr is: the sum of the mass percentages of Mo and Cr is 0.2%.

[0008] The mass percentages of each element in the heat-resistant alloy of the Al-Cu-Mg-Ag-Mn system are as follows:

[0009] Cu: 4.8 - 6.4%, Mg: 0.4 - 0.8%, Ag: 0.2 - 1.2%, Mn: 0.2 - 0.8%, Ti: 0 - 0.2%, Mo: 0 - 0.2% and Cr: 0 - 0.2%, and the sum of the mass percentages of Mo and Cr is 0.2%, with the balance being Al and unavoidable impurities.

[0010] In the Al-Cu-Mg-Ag-Mn series heat-resistant alloy described above, rod-shaped T-Al precipitates after homogenization heat treatment. 20 The Cu2Mn3 dispersion phase presents a B-centered tetragonal structure with lattice parameters Spherical Al7(Cr,Mn) presenting a pseudo-octahedral symmetric rotation axis structure and Al6(Fe,Mn,Mo) dispersion phase with an orthorhombic crystal structure.

[0011] In the Al-Cu-Mg-Ag-Mn series heat-resistant alloy described above, the tensile strength in room-temperature tension is 470 - 510 MPa, the yield strength is 420 - 470 MPa, and the elongation is 9 - 12%; the tensile strength in high-temperature tension at 250 °C is 300 - 340 MPa, the yield strength is 290 - 330 MPa, and the elongation is 15 - 21%.

[0012] In the Al-Cu-Mg-Ag-Mn series heat-resistant alloy described above, when thermally exposed at 200 °C for 100 h and 500 h, the measured best tensile strength is only reduced by 11.8% and 19.4% respectively compared to the room-temperature tensile strength; when thermally exposed at 250 °C for 100 h and 500 h, it is reduced by 22.2% and 29.5%, showing excellent heat resistance.

[0013] The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy of the present invention involves batching and melting according to the mass percentages of each element of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy, followed by double-stage homogenization, deformation processing, solution treatment, and aging treatment to obtain the Al-Cu-Mg-Ag-Mn series heat-resistant alloy described above.

[0014] Specifically, it includes the following steps:

[0015] S1: Melting and casting

[0016] Prepare raw materials according to the mass percentages of each element of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy, and conduct melting, degassing, slag skimming, re-holding and standing, and casting to obtain an ingot;

[0017] S2: Double-stage homogenization

[0018] According to the DSC curve measured by a differential scanning calorimeter, the exothermic peak corresponding to the precipitation of the dispersed phase is at 380 - 420 °C, and the endothermic peak for the melting of the main second phase Al2Cu is at 530 - 560 °C. According to the segregation degree of Cu > Mg > Ag, under the condition of preventing overburning, the segregation degree of Cu determines the homogenization time;

[0019] The ingot is subjected to a two-stage homogenization treatment to obtain a homogenized ingot.

[0020] S3: Deformation processing

[0021] The homogenized ingot is subjected to surface milling, followed by hot rolling, annealing, and cold rolling to obtain rolled sheets with fewer defects such as edge cracking and better quality;

[0022] S4: Solution and aging treatment

[0023] After the rolled sheets are subjected to solution treatment, they are then subjected to aging treatment to obtain an Al-Cu-Mg-Ag-Mn series heat-resistant alloy.

[0024] In the Al-Cu-Mg-Ag-Mn series heat-resistant alloy described in the present invention, Mo and Cr are added. Cr can improve the stress corrosion cracking resistance of the alloy and reduce the tendency of hot cracks at high temperatures; trace amounts of Mo inhibit and eliminate the formation of coarse and insoluble phases containing Fe, improving the formability and fracture toughness of the alloy. Additionally, according to the difference in its equilibrium distribution coefficient (K Mn <1, K Mo >1 and K Cr >1), rod-shaped Al 20 Cu2Mn3 dispersed phase tends to precipitate in the interdendritic region. The addition of Mo and Cr promotes the precipitation of dispersed phases Al6(Fe,Mn,Mo) and Al7Cr, and optimizes the size, morphology, distribution, and number density of the dispersed phases to varying degrees. Uniform and fine dispersed phases not only achieve the effect of dispersion strengthening, but also increase the recrystallization temperature of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy, effectively pinning the grain boundaries and inhibiting the coarsening of grains at high temperatures, which plays a decisive role in the high-temperature stability of the alloy.

[0025] In the S1 described above, the specific process is as follows: First, pure aluminum is melted at a melting temperature of 730 - 750 °C; after heating to 750 - 780 °C, intermediate alloys of Al-Cu, Al-Mn, Al-Mo, Al-Cr, and Al-Ti with higher melting points are added; after full melting, the temperature is lowered to 700 - 720 °C, and then pure Ag and pure Mg are added, and they are pressed into the melt with a gas pressure hood to prevent burning loss. After melting, hexachloroethane is used for degassing and stirring, and it is kept warm and static for 10 min. After the elemental compositions in the melt are evenly distributed, water-cooled copper mold casting is carried out;

[0026] In the aforementioned S2, the double-stage homogenization includes the following steps:

[0027] (1) First, homogenize and anneal the ingot at 400 ± 20 °C for 8 - 16 h. At this temperature, it is beneficial to the precipitation of dispersed phases;

[0028] (2) Then adjust the temperature to 515 ± 15 °C, hold for 12 - 36 h, and then air-cool to room temperature;

[0029] In the aforementioned S3, the deformation processing includes the following steps:

[0030] (1) First, place the homogenized ingot milled to a thickness of 20 - 25 mm at 450 - 480 °C for 1 h to obtain a slab with a thickness of 20 - 25 mm;

[0031] (2) Hot-roll the slab with a thickness of 20 - 25 mm to 5 mm, with a total reduction of 70 - 80% and 6 - 8 passes of processing, to obtain a hot-rolled plate; during hot-rolling, the temperature of the ingot is maintained above 400 °C to ensure good processing performance during deformation;

[0032] Furthermore, use a two-high hot-rolling experimental mill to hot-roll the slab with a thickness of 20 - 25 mm to 5 mm.

[0033] (3) Then anneal the hot-rolled plate at 450 - 480 °C for 1 - 4 h to eliminate work hardening and reduce the deformation resistance, which is beneficial for subsequent cold processing;

[0034] (3) Use a two-high irreversible rolling mill to further cold-roll the annealed slab for 6 - 12 passes, preferably 6 - 8 passes, with a total cold reduction of 30 - 80% and a thickness reduction from 5 mm to 2 mm to obtain the deformed blank.

[0035] In the aforementioned S4, the solution and aging heat treatment include the following steps:

[0036] (1) First, hold the rolled sheet at 500 - 550 °C for 1 - 6 h and water-quench to room temperature to obtain a solution-treated alloy; preferably hold at 500 - 540 °C for 1 - 4 h; more preferably hold at 510 - 530 °C for 1 - 4 h according to hardness and conductivity.

[0037] (2) Then age the solution-treated alloy at 150 - 180 °C. The optimal temperature is 165 °C and the time is 8 - 16 h to obtain the Al-Cu-Mg-Ag-Mn series heat-resistant alloy.

[0038] In step S2, according to the DSC curve, it can be determined that the first stage homogenization at 400±20°C can promote the precipitation of the dispersed phase, and the second stage homogenization controlling the temperature at 515±15°C can dissolve the non-equilibrium eutectic phase, eliminate dendrite segregation, and improve the uniformity of the chemical composition and structure of the alloy.

[0039] In step S3, the hot rolling process in the deformation process can eliminate the as-cast defects such as looseness and porosity, transform the as-cast structure into a deformed structure, and improve the processing performance of the alloy; annealing mainly eliminates residual stress, improves plasticity, reduces deformation resistance, and improves the heterogeneity of the structure and performance after hot rolling; the deformation strengthening caused by cold rolling combined with solid solution and aging heat treatment can further improve the comprehensive performance of the alloy.

[0040] In step S4, the solution treatment is carried out at high temperature and then water quenched, which increases the solubility of strengthening elements Cu and Mg in the matrix, which is conducive to the precipitation of strengthening phases (θ′-Al2Cu and Ω) in the aging stage. In addition, there are a large number of dislocations and substructures in the deformed structure, which gradually swallow up the nuclei in the recovery and recrystallization stages, and form uniform and fine equiaxed crystals as time goes by.

[0041] The Al-Cu-Mg-Ag-Mn heat-resistant alloy prepared by the present invention has high tensile strength and yield strength under room temperature stretching and 250°C high temperature stretching; and when the peak aging sample is exposed to heat for a long time at 200-250°C for 100-500h, the alloy still maintains high strength and has good thermal stability.

[0042] The Al-Cu-Mg-Ag-Mn heat-resistant alloy of the present invention is characterized by using Mo and Cr microalloying, and undergoing the above-mentioned two-stage homogenization, deformation processing, solid solution, and aging heat treatment, and exerting the synergistic effect of Mo and Cr elements on the basis of the best process, thereby reducing the harm of Fe-containing insoluble phase to the mechanical properties of the alloy. Mn <1,K Mo >1 and K Cr >1), the Al6(Fe,Mn,Mo) and Al7(Cr,Mn) dispersed phases formed by the addition of Mo and Cr eliminate the rod-like T-Al 20 The Cu2Mn3 dispersed phase tends to precipitate in the interdendritic region, resulting in precipitation-free zones and sparse precipitation zones, and the size, morphology, distribution and number density of the dispersed phase are optimized to varying degrees.

[0043] The uniform and fine dispersed phase can play a role in dispersion strengthening, improving the mechanical properties of Al-Cu-Mg-Ag-Mn heat-resistant alloys at room temperature and high temperature; it can effectively pin the grain boundaries and inhibit the coarsening of grains at high temperatures, playing a decisive role in the high-temperature stability of the alloy. Description of the Drawings

[0044] Figure 1 : Metallographic diagrams of the S1, S2, S3, and S4 alloys of the present invention after double-stage homogenization and etched with Keller's reagent for 10 - 30 s, from which the influence of the addition of Mo and Cr on the distribution of dispersed phases can be seen.

[0045] Figure 2 : Transmission diagrams of the S1, S2, S3, and S4 alloys of the present invention after double-stage homogenization, showing the changes in the size and morphology of the dispersed phases Al 20 Cu2Mn3, Al6(Fe,Mn,Mo), and Al7(Cr,Mn).

[0046] Figure 3 : Statistics on the average size and number density of rod-shaped T-Al 20 Cu2Mn3, spherical Al7(Cr,Mn), and Al6(Fe,Mn,Mo) in the S1, S2, S3, and S4 alloys of the present invention.

[0047] Figure 4 : EBSD grain boundary diagrams of hot compression of the alloy combinations of the present invention after double-stage homogenization, showing the dynamic recrystallization situation.

[0048] Figure 5 : Room temperature tensile curves of the S1, S2, S3, and S4 alloys of the present invention.

[0049] Figure 6 : High temperature tensile curves of the S1, S2, S3, and S4 alloys of the present invention at 250 °C. Detailed Description of the Invention

[0050] The following non-limiting examples can enable those of ordinary skill in the art to more comprehensively understand the present invention, but do not limit the present invention in any way.

[0051] In the following examples, the test methods are conventional methods unless otherwise specified; the reagents and materials can be obtained from commercial sources unless otherwise specified.

[0052] The present invention will be further described below in conjunction with the examples, but the present invention is not limited to the following examples.

[0053] Example 1:

[0054] Prepare the alloy composition according to Table 1 below. Use a graphite crucible and a water-cooled copper mold for preparation. The raw materials used are pure aluminum, pure magnesium, pure silver, and Al-Cu, Al-Mn, Al-Ti, Al-Mo, Al-Cr master alloys. Melt the aluminum ingot at a melting temperature of 730 - 750 °C, raise the temperature to 750 - 780 °C, then add the Al-Cu, Al-Mn, Al-Mo, Al-Ti, Al-Cr master alloys to fully melt them. Then lower the melt temperature to 700 - 720 °C and add pure magnesium and pure silver. After they melt, degas with hexachloroethane, stir, and keep warm and static for 10 min. After the elemental composition in the melt is evenly distributed, perform water-cooled copper mold casting. Four groups of Al-Cu-Mg-Ag-Mn heat-resistant alloys with different compositions are prepared, and their actual compositions are measured by a spectrometer, as shown in Table 1.

[0055] Table 1. Actual alloy composition

[0056]

[0057] Example 2: Perform double-stage homogenization treatment on the as-cast S1 alloy in Example 1 at 400 °C × 12 h + 515 °C × 24 h, air-cool to room temperature, intercept part of the homogenized specimen, polish it with sandpaper, and immerse it in Keller's reagent for corrosion for 10 - 30 s. Its dispersion phase distribution, size, and number density are as shown in Figure 1 、 Figure 2 and Figure 3 . It can be seen that the rod-shaped T-Al 20 Cu2Mn3 disperses and segregates between dendrites, resulting in a large amount of dispersion-free zone (DFZ) in the grains. The area of the DFZ is 25 - 35%, and the length of T-Al 20 Cu2Mn3 is 275.4 ± 9.5 nm, and the average number density is 3.5 / μm 2 ; The high-throughput double-cone specimen of 500 °C - 0.01 s -1 -60% hot compression is treated at 400 °C × 12 h + 515 °C × 24 h, and its EBSD grain boundary map at the core is as shown in Figure 4 . The proportions of low-angle grain boundaries (LABs), medium-angle grain boundaries (MLABs), and high-angle grain boundaries (HABs) in the grain boundary misorientation difference are 37.52%, 6.25%, and 24.19% respectively. It shows that only large-sized and low-number-density T-Al 20 Cu2Mn3 dispersion phases precipitate during the homogenization stage of the S1 alloy, and the area of the DFZ is large, the dispersion phase distribution is extremely uneven, and the effect of inhibiting recrystallization during the hot compression process is not obvious.

[0058] Case 3: The as-cast S2 alloy in Case 1 was subjected to a two-step homogenization treatment of 400 °C × 12 h + 515 °C × 24 h, air-cooled to room temperature, and a part of the homogenized specimen was intercepted, polished with sandpaper, and etched in Keller's reagent for 10 - 30 s. The distribution, size, and number density of its dispersed phases are as Figure 1 , Figure 2 and Figure 3 shown. It can be seen that the area of the DFZ in the grains decreased significantly, with an area of 15 - 25%, and the length of the T-Al 20 Cu2Mn3 dispersed phase was 156.5 ± 8.1 nm, and the average number density was 7.6 / μm 2 ; In addition to T-Al 20 Cu2Mn3, spherical Al7(Cr,Mn) and Al6(Fe,Mn,Mo) dispersed phases were also precipitated, with a diameter of 127.6 ± 3.2 nm and an average number density of 5.6 / μm 2 . The high-throughput double-cone specimens with 500 °C - 0.01 s -1 -60% hot compression were treated with 400 °C × 12 h + 515 °C × 24 h, and the EBSD grain boundary map of its core is as Figure 4 . The proportions of low-angle grain boundaries (LABs), medium-angle grain boundaries (MLABs), and high-angle grain boundaries (HABs) in the grain boundary misorientation differences were 42.72%, 6.51%, and 18.55% respectively. It shows that adding 0.15 wt% Cr and 0.05 wt% Mo reduced the length of the T-Al 20 Cu2Mn3 dispersed phase by 43.17% on average, the number density increased by 4.1 / μm 2 , and the DFZ also decreased by 10%; In addition, the precipitation of spherical dispersed phases hindered the movement of dislocations and the migration of sub-grain boundaries, inhibiting the transformation from LABs to HABs and increasing the recrystallization temperature of the alloy.

[0059] Case 4: The as-cast S3 alloy in Case 1 was subjected to a two-step homogenization treatment of 400 °C × 12 h + 515 °C × 24 h, air-cooled to room temperature, and a part of the homogenized specimen was intercepted, polished with sandpaper, and etched in Keller's reagent for 10 - 30 s. The distribution, size, and number density of its dispersed phases are as Figure 1 , Figure 2 and Figure 3 shown. It can be seen that the area of the DFZ in the grains increased compared with that of S2, being 20 - 30%, and the length of the T-Al 20 Cu2Mn3 dispersed phase was 172.7 ± 6.7 nm, and the average number density was 5.0 / μm 2 ; The diameters of the Al7(Cr,Mn) and Al6(Fe,Mn,Mo) dispersed phases were 105.4 ± 2.7 nm, and the average number density was 7.7 / μm2 。500 °C - 0.01 s -1 - 60% hot-compressed high-throughput double-cone specimens were treated at 400 °C × 12 h + 515 °C × 24 h, and the EBSD grain boundary map of their core is as Figure 4 。The proportions of low-angle grain boundaries (LABs), medium-angle grain boundaries (MLABs), and high-angle grain boundaries (HABs) in the grain boundary misorientation are 44.96%, 6.02%, and 17.37% respectively. It shows that with the addition of 0.10 wt% Cr and 0.10 wt% Mo, although the DFZ area is 5% more than that of S2 alloy, the length of T-Al 20 Cu2Mn3 increases by 10.35% and the number density decreases by 2.6 / μm 2 ; but the average diameter of the spherical dispersed phase decreases by 22.2 nm and the number density increases by 2.1 / μm 2 , the transformation from LABs to HABs is further inhibited, and the spherical dispersed phase shows a more excellent recrystallization inhibition effect compared to the rod-shaped dispersed phase.

[0060] Example 5: The as-cast S4 alloy in Example 1 was subjected to a two-stage homogenization treatment at 400 °C × 12 h + 515 °C × 24 h, air-cooled to room temperature, and part of the homogenized specimens were intercepted, polished with sandpaper, and placed in Keller's reagent for corrosion for 10 - 30 s. The distribution, size, and number density of its dispersed phases are as Figure 1 、 Figure 2 and Figure 3 shown. It can be seen that as the Mo percentage content increases to 0.15%, the precipitation of the Al6(Fe,Mn,Mo) dispersed phase into the dendrite core region basically eliminates the DFZ and the distribution becomes more uniform. T-Al 20 The size of the Cu2Mn3 dispersed phase is 183.5 ± 7.5 nm, and the average number density is 4.5 / μm 2 ; the size of the Al7(Cr,Mn) and Al6(Fe,Mn,Mo) dispersed phases is 86.2 ± 4.6 nm, and the average number density is 10.3 / μm 2 。500 °C - 0.01 s -1 - 60% hot-compressed high-throughput double-cone specimens were treated at 400 °C × 12 h + 515 °C × 24 h, and the EBSD grain boundary map of their core is as Figure 4 。The proportions of low-angle grain boundaries (LABs), medium-angle grain boundaries (MLABs), and high-angle grain boundaries (HABs) in the grain boundary misorientation are 44.83%, 5.11%, and 14.75% respectively. It shows that with the addition of 0.05 wt% Cr and 0.15 wt% Mo, the uniform distribution of the dispersed phase is promoted, the DFZ is completely eliminated, and the spherical dispersed phase has a small size and the largest number density, effectively hindering the dislocation movement and sub-grain boundary migration, and the recrystallization inhibition effect is the most obvious.

[0061] Implementation Case 6: The 4 groups of Al-Cu-Mg-Ag-Mn alloys in Implementation Cases 2-5 that have undergone double-stage homogenization treatment at 400°C × 12 h + 515°C × 24 h are subjected to deformation processing. The specific process is as follows: hot rolling is carried out at 470°C, the total deformation is 80%, the number of passes is 6, the hot-rolled plate is annealed at 475°C for 2 h to eliminate residual stress, and then cold rolling is carried out after air cooling. The total deformation is 60%, and the number of passes is 12; then solution aging treatment is carried out. The specific process is as follows: the cold-rolled plate is solution-treated at 525°C, water quenched immediately after holding for 1 h, and then aged at 165°C for 8 h. Room temperature tensile and high-temperature tensile tests at 250°C are carried out on the 4 groups of Al-Cu-Mg-Ag-Mn, and the tensile curves are respectively as Figure 5 and Figure 6 shown. As can be seen from Figure 5 , the addition of Cr and Mo improves the room temperature tensile strength of the alloy to varying degrees. Among them, the strength of alloy S4 reaches the maximum of 505.7 MPa, the yield strength is 464.9 MPa, and the elongation is 9.7%. Compared with the tensile strength and yield strength of alloy S1, they are increased by 27.8 MPa and 35.3 MPa respectively, and the elongation changes little. As can be seen from Figure 6 , the addition of Cr and Mo improves the high-temperature tensile strength of the alloy to varying degrees. Among them, the high-temperature strength of alloy S4 reaches the maximum of 338.7 MPa, the yield strength is 329.6 MPa, and the elongation is 19.0%. Compared with the tensile strength and yield strength of alloy S1, they are increased by 28.0 MPa and 28.1 MPa respectively.

[0062] Implementation Case 7: A thermal exposure test is carried out on the S1 thin plate in Implementation Case 6 that has undergone double-stage homogenization at 400°C × 12 h + 515°C × 24 h, hot rolling and cold rolling, and solution at 525°C for 1 h and aging at 165°C for 8 h. The specific steps are as follows: thermal exposure treatments are carried out at 200°C and 250°C for 100 h and 500 h, and its mechanical properties are tested. The results are shown in Table 2.

[0063] Implementation Case 8: A thermal exposure test is carried out on the S2 thin plate in Implementation Case 4 that has undergone double-stage homogenization at 400°C × 12 h + 515°C × 24 h, hot rolling and cold rolling, and solution at 525°C for 1 h and aging at 165°C for 8 h. The specific steps are as follows: thermal exposure treatments are carried out at 200°C and 250°C for 100 h and 500 h, and its mechanical properties are tested. The results are shown in Table 3.

[0064] Implementation Case 9: A heat exposure test was conducted on the S3 thin plate that had undergone double-stage homogenization at 400°C × 12 h + 515°C × 24 h, hot rolling, cold rolling, solution treatment at 525°C for 1 h, and aging at 165°C for 8 h in Implementation Case 4. The specific steps were as follows: heat exposure treatments were carried out at 200°C and 250°C for 100 h and 500 h, respectively, and mechanical property tests were performed on it. The results are shown in Table 4.

[0065] Implementation Case 10: A heat exposure test was conducted on the S4 thin plate that had undergone double-stage homogenization at 400°C × 12 h + 515°C × 24 h, hot rolling, cold rolling, solution treatment at 525°C for 1 h, and aging at 165°C for 8 h in Implementation Case 4. The specific steps were as follows: heat exposure treatments were carried out at 200°C and 250°C for 100 h and 500 h, respectively, and mechanical property tests were performed on it. The results are shown in Table 5.

[0066] Table 2

[0067]

[0068] Table 3

[0069]

[0070] Table 4

[0071]

[0072]

[0073] Table 5

[0074]

Claims

1. A heat-resistant alloy of the Al-Cu-Mg-Ag-Mn system, characterized in that: The Al-Cu-Mg-Ag-Mn heat-resistant alloy comprises an Al-Cu-Mg-Ag-Mn matrix and trace amounts of Mo and Cr, and the relationship between Mo and Cr is that the sum of the mass percentages of Mo and Cr is 0.2%; The chemical components contained in the said Al-Cu-Mg-Ag-Mn matrix are by mass percentage: Cu: 4.8 - 6.4%, Mg: 0.4 - 0.8%, Ag: 0.2 - 1.2%, Mn: 0.2 - 0.8%, Ti: 0 - 0.04%, and the balance is Al and unavoidable impurities.

2. The heat-resistant alloy of the Al-Cu-Mg-Ag-Mn system according to claim 1, characterized in that: The described Al-Cu-Mg-Ag-Mn heat-resistant alloy precipitates rod-shaped T-Al 20 Cu2Mn3 dispersoid phases, presenting a B-centered tetragonal structure with lattice parameters Spherical Al7(Cr,Mn) with a pseudo-octahedral symmetric rotation axis structure and Al6(Fe,Mn,Mo) dispersoid phases with an orthorhombic crystal structure.

3. The Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 1, characterized in that: The tensile strength of the said Al-Cu-Mg-Ag-Mn heat-resistant alloy in room-temperature tension is 470 - 510 MPa, the yield strength is 420 - 470 MPa, and the elongation is 9 - 12%; in 250°C high-temperature tension, the tensile strength is 300 - 340 MPa, the yield strength is 290 - 330 MPa, and the elongation is 15 - 21%; For the said Al-Cu-Mg-Ag-Mn heat-resistant alloy, when thermally exposed at 200°C for 100 h and 500 h, the measured best tensile strength is only reduced by 11.8% and 19.4% respectively compared with the room-temperature tensile strength; when thermally exposed at 250°C for 100 h and 500 h, it is reduced by 22.2% and 29.5%, showing excellent heat-resistant performance.

4. The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 1, characterized in that: According to the mass percentages of each element of the Al-Cu-Mg-Ag-Mn heat-resistant alloy for batching and melting and casting, and then carrying out double-stage homogenization, deformation processing, solution treatment and aging treatment to obtain the said Al-Cu-Mg-Ag-Mn heat-resistant alloy.

5. The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 4, characterized in that: The said melting and casting is as follows: First, melt pure Al, control the temperature at 730 - 750°C, then add Al-Cu, Al-Mn, Al-Ti, Al-Mo and Al-Cr master alloys, control the temperature at 750 - 780°C, after it melts, then control the melting temperature at 700 - 720°C and add pure Ag and pure Mg, after it melts sufficiently, degas with hexachloroethane, skim the slag, stir, keep warm and stand still, and carry out water-cooled copper mold casting after the melt element distribution is uniform.

6. The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 4, characterized in that: The said double-stage homogenization process is: The first stage: keep warm at 400 ± 20°C for 8 - 16 h; the second stage: keep warm at 515 ± 15°C for 12 - 36 h; then air-cool to room temperature.

7. The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 4, characterized in that: The said deformation processing is: Carry out hot rolling and cold rolling. Among them, the hot rolling temperature is controlled at 450 - 480°C, the total reduction is 70 - 80%, and the number of processing passes is 6 - 8 times; the intermediate annealing temperature is 450 - 480°C, and the time is 1 - 4 h; the total cold rolling reduction is 30 - 80%, and the number of processing passes is 6 - 12 times.

8. The preparation method of the Al-Cu-Mg-Ag-Mn series heat-resistant alloy according to claim 4, characterized in that: The temperature of the said solution treatment is controlled at 500 - 550°C, the holding time is 1 - 6 h, and it is immediately water quenched after taking out; The temperature of the said aging treatment is controlled at 150 - 180°C, and the holding time is 8 - 16 h.

Images