A high-strength aluminum alloy having excellent fatigue damage resistance and a method for manufacturing the same

Abstract

The application provides a high-strength aluminum alloy with excellent anti-fatigue damage performance, and belongs to the technical field of non-ferrous metal material preparation. The high-strength aluminum alloy contains the following components in mass fraction: Cu: 4.5-5.5%, Mg: 1.0-1.4%, Ag: 0.2-1%, Mn: 0.5-1%, Ti: 0.06-0.10%, Fe: less than 0.06%, Si: less than 0.06%, and Zn: less than 0.1%; the balance is Al, and 3<=Cu / Mg<=4. The microstructure of the high-strength aluminum alloy contains Goss+Cube or Goss+P grain bands, the grain length-diameter in the grain bands is parallel to the TD direction, the grain boundary is perpendicular to the RD direction, and equiaxed grains are distributed around the grain bands. The high-strength aluminum alloy has high strength and toughness and excellent anti-fatigue damage performance, the material component design is reasonable, the preparation process is simple, and the industrial production cost is low. The application further provides a preparation method of the high-strength aluminum alloy with excellent anti-fatigue damage performance.

Description

TECHNICAL FIELD

[0001] The present application belongs to the technical field of non-ferrous metal material preparation, and particularly relates to a high-strength aluminum alloy with excellent fatigue damage resistance and a preparation method thereof. BACKGROUND

[0002] 2000 series aluminum alloys have moderate strength and good fatigue resistance, and are main candidate materials for aircraft structural components and aircraft skin. Typical alloys are 2x24 series alloys developed by the Alcoa Company of the United States, such as 2524 alloy, which is an improved version based on 2024 alloy. By controlling Cu, Mg and reducing Fe and Si content, the fracture toughness and fatigue performance are effectively improved, but the alloy strength is reduced. The 2029 alloy developed in recent years also has relatively excellent fracture toughness and fatigue resistance, but the alloy strength is also relatively low. With the development of the aviation industry, Al-Cu-Mg-based alloys are developing towards high strength, excellent fatigue damage resistance and low preparation cost.

[0003] The internal factors affecting fatigue performance mainly include grain size, second phase particles, grain boundary structure and texture. The authorized patent technologies for improving fatigue performance through texture component design are mainly CN103045976A, CN103526140A, CN10358997A, CN108103373B and CN108504915B published by the inventors in recent years. The fatigue resistance of the Al-Cu-Mg alloys disclosed in these technologies is relatively excellent, but the tensile strength is generally low (<480 MPa). The technology disclosed in patent CN10358997A can prepare an alloy with a tensile strength exceeding 480 MPa, but the preparation process is complicated and the microstructure stability is difficult to guarantee. In addition, patents CN105506521 and CN108149172B disclose a processing method for fatigue-resistant aluminum alloy sheets with brass texture or fine-grained cubic texture, and the ultimate tensile strength of the alloy is only about 400 MPa. In the currently disclosed patent technologies, the yield strength ratio of the alloy is relatively high, and when the yield strength ratio of the alloy increases, the plastic damage space decreases, which means that the dislocation slip will be inhibited, increasing the stress concentration in the fatigue process and further increasing the tendency of micro-crack formation, thereby reducing the fatigue damage resistance. It can be seen that it is very difficult to make the alloy have high strength and toughness and excellent fatigue damage resistance, and the existing disclosed patent technologies cannot solve the problem that the Al-Cu-Mg-based alloy has an ultimate tensile strength exceeding 500 MPa while still having excellent fatigue damage resistance.

[0004] In summary, the 2000 series aluminum alloys disclosed in the prior art are very difficult to have high strength and toughness, excellent fatigue damage resistance and low preparation cost. With the development of the aviation industry, the performance indicators of the alloys in the currently disclosed patent technologies have been difficult to meet the design requirements of future aircraft. SUMMARY

[0005] In order to solve the problem that the aluminum alloy is difficult to have high strength and toughness and excellent fatigue damage resistance, the application provides a high-strength aluminum alloy with excellent fatigue damage resistance, which has high strength and toughness and excellent fatigue damage resistance, has reasonable material component design, simple preparation process and low industrial production cost.

[0006] The application also provides a preparation method of the high-strength aluminum alloy with excellent fatigue damage resistance.

[0007] The application achieves the above-mentioned technical effects through the following technical scheme.

[0008] The application provides a high-strength aluminum alloy with excellent fatigue damage resistance, and the chemical components of the high-strength aluminum alloy are as follows in terms of mass fraction:

[0009] Cu: 4.5-5.5%, Mg: 1.0-1.4%, Ag: 0.2-1%, Mn: 0.5-1%, Ti: 0.06-0.10%, Fe <0.06%, Si <0.06%, Zn <0.1%; the balance is Al, 3≤Cu / Mg≤4;

[0010] The microstructure of the high-strength aluminum alloy contains Goss+Cube or Goss+P grain bands, the grain length-diameter in the grain bands is parallel to the TD direction, the grain boundary is perpendicular to the RD direction, and equiaxed grains are distributed around the grain bands.

[0011] Preferably, the yield strength of the high-strength aluminum alloy is 300-350 MPa, the ultimate tensile strength is 495-520 MPa, and the elongation after fracture is 18-30%.

[0012] The fatigue damage resistance index of the high-strength aluminum alloy in the TL direction is as follows: ΔK=20 MPa*m 1 / 2 , da / dN <5×10 -4 mm / cycle; ΔK=30 MPa*m 1 / 2 , da / dN <3×10 -3 mm / cycle; ΔK=33 MPa*m 1 / 2 , da / dN <4×10 -3 mm / cycle, and the instability ΔK can reach 39 MPa*m 1 / 2 .

[0013] Based on the same inventive concept, the application provides a preparation method of the high-strength aluminum alloy with excellent fatigue damage resistance, and the preparation method comprises the following steps:

[0014] A cast ingot is obtained through a melting and casting process.

[0015] homogenizing the ingot to obtain a homogenized ingot;

[0016] hot-rolling the homogenized ingot to obtain a hot-rolled plate;

[0017] or, hot-rolling the homogenized ingot, annealing, cold-rolling or directly cold-rolling to obtain a cold-rolled plate;

[0018] isothermal annealing the hot-rolled plate or the cold-rolled plate to obtain an annealed aluminum alloy plate;

[0019] high-temperature short-time solid solution treatment, water quenching and aging treatment to obtain a high-strength aluminum alloy;

[0020] wherein the chemical composition of the ingot is as follows in mass fraction:

[0021] Cu: 4.5-5.5%, Mg: 1.0-1.4%, Ag: 0.2-1%, Mn: 0.5-1%, Ti: 0.06-0.10%, Fe <0.06%, Si <0.06%, Zn <0.1%; the balance being Al, 3≤Cu / Mg≤4.

[0022] Preferably, the homogenizing the ingot to obtain a homogenized ingot comprises:

[0023] homogenizing the ingot at 465-490℃ for 24-48h;

[0024] after the first homogenization, heating to 490-505℃ for the second homogenization for 12-48h to obtain the homogenized ingot.

[0025] Preferably, the hot-rolling the homogenized ingot to obtain a hot-rolled plate comprises:

[0026] hot-rolling the homogenized ingot at 390-490℃ with a total deformation of more than 85% to obtain a hot-rolled plate with a thickness of 2-6mm.

[0027] Preferably, the hot-rolling the homogenized ingot, annealing, cold-rolling or directly cold-rolling to obtain a cold-rolled plate comprises:

[0028] homogenizing treatment ingot is subjected to large deformation hot rolling at 390-490 DEG C, the total deformation amount is greater than 85%, and a hot-rolled plate with a thickness of 2-6 mm is obtained, then the hot-rolled plate is directly subjected to cold rolling, the cold rolling deformation amount is less than or equal to 50%, and a cold-rolled plate is obtained;

[0029] or, the homogenizing treatment ingot is subjected to large deformation hot rolling at 390-490 DEG C, the total deformation amount is greater than 85%, and a hot-rolled plate with a thickness of 2-6 mm is obtained, the hot-rolled plate is subjected to annealing treatment, the annealing temperature is 260 DEG C-330 DEG C, the annealing time is 2-4 h, then the annealed plate is subjected to cold rolling, the cold rolling deformation amount is less than or equal to 50%, and a cold-rolled plate is obtained.

[0030] Preferably, the hot-rolled plate or the cold-rolled plate is subjected to isothermal annealing, and an annealed aluminum alloy plate is obtained, and the method specifically comprises the following steps:

[0031] the hot-rolled plate or the cold-rolled plate is subjected to isothermal annealing, and an annealed aluminum alloy plate is obtained;

[0032] wherein, when the large deformation hot rolling temperature is lower than 450 DEG C, the isothermal annealing temperature is controlled to be 320 DEG C-370 DEG C, and the annealing time is 2-5 h, when the large deformation hot rolling temperature is higher than 450 DEG C, the isothermal annealing temperature is controlled to be 340 DEG C-420 DEG C, and the annealing time is 3-6 h.

[0033] Preferably, the annealed aluminum alloy plate is subjected to high-temperature short-time solid solution treatment, water quenching and aging treatment, and a high-strength aluminum alloy is obtained, and the method specifically comprises the following steps:

[0034] the annealed aluminum alloy plate is subjected to high-temperature short-time solid solution treatment, the solid solution treatment temperature is 475-505 DEG C, and the time is 3-30 min, then water quenching is carried out by taking water as a medium, the quenching transfer time is less than or equal to 15 s, and then aging treatment is carried out, and a high-strength aluminum alloy is obtained;

[0035] wherein, when the aging treatment adopts natural aging system, the natural aging time is greater than or equal to 240 h, when the aging treatment adopts artificial aging system, the artificial aging temperature is 140 DEG C-180 DEG C, and the aging time is 1-5 h.

[0036] Based on the same inventive concept, the application further provides an application of the high-strength aluminum alloy with excellent fatigue damage resistance in preparation of an aviation or aerospace vehicle structural member.

[0037] The one or more technical solutions in the embodiments of the application have at least the following technical effects or advantages:

[0038] 1.The high-strength aluminum alloy with excellent anti-fatigue damage performance, which solves the problem that the existing Al-Cu-Mg-based alloy cannot simultaneously achieve high strength and excellent anti-fatigue damage performance, and meets the low-cost preparation concept, the high-strength aluminum alloy prepared by the application has a tensile strength of 500 MPa, an elongation of 18-30%, and excellent strength and toughness, the fatigue crack propagation rate of the aged alloy plate in the TL direction is: ΔK=20 MPa*m 1 / 2 , da / dN<5×10 - 4 mm / cycle; ΔK=30 MPa*m 1 / 2 , da / dN<3×10 -3 mm / cycle; ΔK=33 MPa*m 1 / 2 , da / dN<4×10 - 3 mm / cycle, and the instability ΔK can reach 39 MPa*m 1 / 2 The above-mentioned high-strength aluminum alloy has excellent anti-fatigue damage performance, high strength and toughness, and excellent anti-fatigue damage performance, the component design of the high-strength aluminum alloy material is reasonable, the preparation process is simple, and the application meets the design requirements of current and future aircraft, and is suitable for large-scale industrial production and application.

[0039] 2.The preparation method of the high-strength aluminum alloy with excellent anti-fatigue damage performance, which can control the texture and microstructure of the aged Al-Cu-Mg-based alloy through deformation and micro-alloying, so that the high-strength aluminum alloy can have an ultimate tensile strength of more than 500 MPa and excellent anti-fatigue damage performance, the high-strength aluminum alloy prepared by the method can maintain high strength and toughness and excellent anti-fatigue damage performance in the aged state even after different deformations and heat treatments, and the application solves the problem that the current Al-Cu-Mg-based alloy cannot simultaneously achieve high strength and toughness, excellent anti-fatigue damage, and low-cost preparation, and meets the design index of current and future aircraft. BRIEF DESCRIPTION OF DRAWINGS

[0040] In order to more clearly illustrate the technical solutions in the embodiments of the application, the drawings needed in the embodiment description will be briefly introduced. Obviously, the drawings in the following description are some embodiments of the application, and other drawings can be obtained by those skilled in the art without creative labor.

[0041] Figure 1 The orientation distribution function ODF graph of the hot-rolled state plate of the alloy of Example 1;

[0042] Figure 2Orientation distribution function ODF plot for the as- naturally aged sheet of Example 1;

[0043] Figure 3 Fatigue crack growth rate curve for the Example 1 alloy;

[0044] Figure 4 TEM microstructure for the as-naturally aged Example 1 alloy;

[0045] Figure 5 Orientation distribution function ODF plot for the hot rolled sheet of Example 2 alloy;

[0046] Figure 6 Orientation distribution function ODF plot for the as-naturally aged sheet of Example 2;

[0047] Figure 7 Fatigue crack growth rate curve for the Example 2 alloy;

[0048] Figure 8 TEM microstructure for the as-naturally aged Example 2 alloy;

[0049] Figure 9 Fatigue crack growth rate curve for the Example 3 alloy;

[0050] Figure 10 Orientation distribution function ODF plot for the as-hot rolled and directly cold rolled sheet of Example 4 alloy;

[0051] Figure 11 Orientation distribution function ODF plot for the as-naturally aged sheet of Example 4;

[0052] Figure 12 Fatigue crack growth rate curve and tensile and fatigue performance indices for the Example 4 alloy;

[0053] Figure 13 TEM microstructure for the as-naturally aged Example 4 alloy;

[0054] Figure 14 Schematic of grain structure distribution for the as-naturally aged sheet of Example 4;

[0055] Figure 15 Orientation distribution function ODF plot for the as-naturally aged sheet of Example 5;

[0056] Figure 16 Fatigue crack growth rate curve for the Example 5 alloy;

[0057] Figure 17 TEM microstructure for the as-naturally aged Example 6 alloy;

[0058] Figure 18 Fatigue crack growth rate curve for the Example 6 alloy;

[0059] Figure 19 TEM microstructure of the Example 6 alloy in artificially aged condition;

[0060] Figure 20 Fatigue crack growth rate curve for the Example 7 alloy;

[0061] Figure 21 TEM microstructure of the Example 7 alloy in artificially aged condition;

[0062] Figure 22 Orientational distribution function (ODF) plot for the Comparative Example 1 alloy hot rolled sheet;

[0063] Figure 23 Orientational distribution function (ODF) plot for the Comparative Example 1 naturally aged sheet;

[0064] Figure 24 Fatigue crack growth rate curve for the Comparative Example 1 alloy;

[0065] Figure 25 TEM microstructure of the Comparative Example 1 alloy in naturally aged condition;

[0066] Figure 26 Orientational distribution function (ODF) plot for the Comparative Example 2 alloy hot rolled sheet;

[0067] Figure 27 Orientational distribution function (ODF) plot for the Comparative Example 2 naturally aged sheet;

[0068] Figure 28 Fatigue crack growth rate curve for the Comparative Example 2 alloy;

[0069] Figure 29 TEM microstructure of the Comparative Example 2 alloy in naturally aged condition;

[0070] Figure 30 Fatigue crack growth rate curve for the Comparative Example 3 alloy;

[0071] Figure 31 TEM microstructure of the Comparative Example 3 alloy in artificially aged condition;

[0072] Figure 32 Orientational distribution function (ODF) plot for the Comparative Example 4 alloy hot rolled sheet;

[0073] Figure 33 Orientational distribution function (ODF) plot for the Comparative Example 4 naturally aged sheet;

[0074] Figure 34 Fatigue crack growth rate curve for the Comparative Example 4 alloy;

[0075] Figure 35TEM microstructure of the alloy of Comparative Example 4 in the natural aging state. DETAILED DESCRIPTION

[0076] The advantages and various effects of the present application will be more clearly apparent from the following specific embodiments and examples. It should be understood by those skilled in the art that the specific embodiments and examples are used to illustrate the present application, not to limit the present application.

[0077] Throughout the specification, unless otherwise specifically indicated, the terms used herein are to be understood in the manner as commonly used in the art. Therefore, unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this application belongs. In case of conflict, the present specification prevails.

[0078] Unless otherwise specifically indicated, various raw materials, reagents, instruments and equipment used in the present application can be purchased from the market or can be prepared by existing methods.

[0079] The technical principle of the present application is as follows:

[0080] The present application provides a preparation method of a high-strength aluminum alloy with excellent anti-fatigue damage performance, comprising:

[0081] S1. obtaining an ingot by a melting and casting process;

[0082] S2. performing primary homogenization treatment on the ingot, and then performing secondary homogenization treatment after temperature rising to obtain a homogenization treated ingot;

[0083] S3. performing large deformation hot rolling on the homogenization treated ingot to obtain a hot-rolled plate;

[0084] or, performing large deformation hot rolling on the homogenization treated ingot, and then performing cold rolling after annealing or directly performing cold rolling to obtain a cold-rolled plate;

[0085] S4. performing isothermal annealing on the hot-rolled plate or the cold-rolled plate to obtain an annealed aluminum alloy plate;

[0086] S5. performing high-temperature short-time solid solution treatment, water quenching and aging treatment on the annealed aluminum alloy plate to obtain a high-strength aluminum alloy;

[0087] wherein, in terms of mass fraction, the chemical composition of the ingot is as follows:

[0088] Cu: 4.5-5.5%, Mg: 1.0-1.4%, Ag: 0.2-1%, Mn: 0.5-1%, Ti: 0.06-0.10%, Fe <0.06%, Si <0.06%, Zn <0.1%; the balance is Al, 3≤Cu / Mg≤4.

[0089] In the present application, the ingot adopts the beneficial effects of the above chemical elements and the ratio as follows:

[0090] Cu: 4.5-5.5%, in this range, the solid solution strengthening provided by Cu element and the cluster strengthening provided by the combination of Cu and Mg element can be maximized. Below this mass fraction range, the solid solution strengthening provided by Cu element is significantly weakened. Above this mass fraction range, Cu element will form excessive coarse Al2Cu phase with Al, which is difficult to be completely dissolved into the matrix after homogenization treatment or subsequent solid solution treatment due to reaching the solid solution limit of Cu in Al. The coarse Al2Cu phase not only weakens the strength, but also significantly hinders the dislocation reciprocating slip in the fatigue process, greatly increasing the fatigue crack propagation rate.

[0091] Mg: 1.0-1.4%, in this range, the solid solution strengthening provided by Mg element and the cluster strengthening effect provided by the combination of Mg and Cu element in the aging state can be maximized. The design of Mg element is matched with Cu element. When the mass fraction of Cu element is determined, if the content of Mg element is below this range, the solid solution strengthening and cluster strengthening will be weakened. If above this range, the content of Al2CuMg phase formed by Mg, Cu and Al in the as-cast state is high, which greatly increases the difficulty of homogenization and the risk of overburning.

[0092] Ag: 0.2-1%, the addition of Ag can provide certain solid solution strengthening and contribute to precipitation strengthening. In this mass fraction range, the Ag element in the natural aging state can significantly reduce the dislocation content formed by the collapse of supersaturated vacancies, which is beneficial to the dislocation reciprocating slip in the fatigue process and increases the fatigue crack propagation resistance; on the other hand, it greatly promotes the precipitation of Cu-Mg atomic clusters during artificial aging, shortens the aging time, and reduces the material preparation cost. The precipitated high-density and appropriately sized Cu-Mg atomic clusters not only provide precipitation strengthening, but also are not easy to cause cyclic softening effect, ensuring high strength while the fatigue performance is also excellent. Below this mass fraction range, the ability to inhibit dislocation formation in the natural aging state is poor, and the artificial aging state cannot effectively promote the formation of clusters. Above this mass fraction range, the material preparation cost will be significantly increased.

[0093] Mn: 0.5-1%, the addition of Mn element is to produce Al 20 Cu2Mn3 phase, which mainly precipitates during homogenization. It can provide dispersion strengthening, and Al 20 Cu2Mn3 phase can hinder the grain boundary migration during recrystallization process, contributing to grain refinement and fine-grain strengthening. Below this mass fraction range, the precipitated Al 20The Cu2Mn3 phase has a low content and poor dispersion strengthening effect, and the subsequent grain refinement effect is not obvious. Higher than the mass fraction range, a part of Mn elements will form a high melting point FeMn phase with impurity Fe elements, eliminating the adverse effect of Fe on the mechanical properties of the alloy.

[0094] Ti: 0.06-0.10%, the addition of Ti element is mainly to provide solid solution strengthening, and secondly to produce certain grain refinement effect. In this range, better strengthening effect can be ensured, and the content of coarse phases formed by Ti, Al and Si can be reduced, and the material preparation cost can be reduced.

[0095] Fe: <0.06%, Si: <0.06%, Fe and Si are impurity elements, and the content of <0.06% is mainly to reduce the content of high melting point coarse brittle phases in the matrix and the consumption of Mg elements by Si, thereby improving the mechanical and fatigue properties of the alloy.

[0096] Zn: <0.1%, Zn is an impurity element introduced in the smelting process, and the control of Zn <0.1% can effectively reduce the consumption of Mg elements by Zn. If the content of Zn is too high, Zn atoms will pin the grain boundary migration during the recrystallization process, inhibiting the formation of strong Goss texture.

[0097] The advantage of 3≤Cu / Mg≤4 is that in this Cu / Mg ratio range, the number density, size and distribution of coarse Al2CuMg and Al2Cu second phases precipitated in the as-cast alloy can be effectively controlled, and these phases are easy to dissolve into the matrix after homogenization treatment. The reduction of the content of these coarse second phases means that the solid solution strengthening and precipitation strengthening can be maximized, and the fatigue crack source is also reduced. On the other hand, the reduction of the content of these coarse phases is beneficial to the formation of high-strength Goss and P textures after annealing. Therefore, controlling 3≤Cu / Mg≤4 is beneficial to simultaneously improving the strength and fatigue performance of the alloy.

[0098] Preferably, step S2 specifically comprises:

[0099] The ingot is subjected to primary homogenization treatment at 465-490°C for 24-48h;

[0100] After the primary homogenization treatment, the temperature is raised to 490-505°C for secondary homogenization treatment, and the homogenization time is 12-48h, to obtain a homogenization treated ingot.

[0101] In the present application, the primary homogenization treatment temperature is 465-490°C, which has the advantage that in this temperature range, the Al2CuMg phase in the as-cast alloy can be dissolved into the matrix as much as possible, and the phases are not overburned. Lower than this temperature range, the Al2CuMg phase is difficult to be completely eliminated, and higher than this temperature range, the alloy is overburned.

[0102] In the present application, the advantage of the two-stage homogenization treatment temperature of 490-505℃ is that it can ensure that the Al2Cu phase in the as-cast alloy is dissolved into the matrix as much as possible. Below this temperature range, the Al2Cu phase is difficult to be completely eliminated, and above this temperature range, the alloy is overburned.

[0103] In the present application, the purpose of the two-stage isothermal homogenization treatment (i.e. the heat treatment is at the same temperature during the homogenization process, and the variable is time) is to maximize the dissolution of the coarse Al2CuMg and Al2Cu phases that are supersaturated in the alloy into the matrix.

[0104] Preferably, step S3 specifically comprises:

[0105] The homogenization treated ingot is subjected to large deformation hot rolling at 390-490℃, and the total deformation amount is greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6mm.

[0106] Alternatively, the homogenization treated ingot is subjected to large deformation hot rolling at 390-490℃, and the total deformation amount is greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6mm, and then the hot-rolled plate is directly subjected to cold rolling with a deformation amount of ≤50%, to obtain a cold-rolled plate.

[0107] In the present application, the large deformation hot rolling temperature is 390-490℃, at which the plate forming effect is best, and high-strength Copper and Brass deformation textures can be obtained. If the hot rolling temperature is too low, the plate will crack seriously, and if the hot rolling temperature is too high, the risk of overburning and cracking of the plate will increase, and in addition, the expected texture components cannot be obtained. The total deformation amount of greater than 85% can obtain high-strength Copper and Brass textures.

[0108] Further, the hot-rolled plate is subjected to annealing treatment before cold rolling, and the annealing temperature is 260-330℃ and the annealing time is 2-4h.

[0109] In the present application, the purpose of annealing before cold rolling is to promote the enhancement of Brass texture of the hot-rolled plate, and then promote the formation of recrystallized Goss, and promote the formation of Goss grain bands (the long diameter of the grains inside the grain band is parallel to the TD direction).

[0110] Preferably, step S4 specifically comprises:

[0111] The hot-rolled plate or the cold-rolled plate is subjected to isothermal annealing to obtain an annealed aluminum alloy plate.

[0112] When the large deformation hot rolling temperature is lower than 450 DEG C, the isothermal annealing temperature is controlled in 320 DEG C-370 DEG C, and the annealing time is 2-5h; when the large deformation hot rolling temperature is higher than 450 DEG C, the isothermal annealing temperature is controlled in 340 DEG C-420 DEG C, and the annealing time is 3-6h.

[0113] When the large deformation hot rolling temperature is lower than 450 DEG C, the isothermal annealing temperature is controlled in 320 DEG C-370 DEG C, and the annealing time is 2-5h in the application, the benefit is that high-strength Goss and P textures can be obtained; when the large deformation hot rolling temperature is higher than 450 DEG C, the isothermal annealing temperature is controlled in 340 DEG C-420 DEG C, and the annealing time is 3-6h, the benefit is that high-strength Goss and Cube textures can be obtained at the temperature and time.

[0114] Preferably, the step S5 specifically comprises:

[0115] The annealed aluminum alloy plate is subjected to high-temperature short-time solid solution treatment at a temperature of 475-505 DEG C for 3-30min, and then is quenched in water with a quenching transfer time of less than or equal to 15s, and is subjected to aging treatment to obtain a high-strength aluminum alloy.

[0116] When the aging treatment adopts a natural aging system, the natural aging time is greater than or equal to 240h; when the aging treatment adopts an artificial aging system, the artificial aging temperature is 140 DEG C-180 DEG C, and the aging time is 1-5h.

[0117] In the application, the solid solution treatment temperature is 475-505 DEG C, and the time is 3-30min, the benefit is that in the temperature and time range, the strength of Goss, P and Cube textures formed by isothermal annealing is not reduced, and fine and dispersed S phase particles in the matrix can be fully dissolved into the matrix, so that the solid solution and precipitation strengthening effects are ensured, the quenching is performed in water with a quenching transfer time of less than or equal to 15s, the benefit is that the alloy is in a high solid solubility, which is beneficial to promoting the precipitation of Cu-Mg atom clusters in the aging process.

[0118] In the application, when the aging treatment adopts a natural aging system, the natural aging time is greater than or equal to 240h, the benefit is that the solid solution atoms can be precipitated to generate high-density Cu-Mg atom clusters, and the alloy is in a relatively stable state; when the aging treatment adopts an artificial aging system, the artificial aging temperature is 140 DEG C-180 DEG C, and the aging time is 1-5h, the benefit is that in the temperature and time range, Cu-Mg atom clusters with high number density and appropriate size can be formed.

[0119] The yield strength of the aged aluminum alloy plate prepared by the above technical route is 300-350MPa, the tensile strength is 495-520MPa, and the elongation after fracture is 18-30%. The anti-fatigue damage index of the aged alloy in the TL direction is: Delta K = 20MPa*m1 / 2 da / dN < 5 x 10 -4 mm / cycle; AK = 30 MPa*m 1 / 2 da / dN < 3 x 10 -3 mm / cycle; AK = 33 MPa*m 1 / 2 da / dN < 4 x 10 -3 mm / cycle, unstable AK up to 39 MPa*m 1 / 2 The aged alloy is controlled by strong Goss+Cube or Goss+P texture, the combined strength of the two textures > 10, the volume fraction > 20%, and the combined Goss and P texture component strength > 7. A certain number of Goss+Cube or Goss+P grain bands are contained in the matrix, the grain length of the grain bands inside is parallel to the TD direction, and the grain boundary is perpendicular to the RD direction. Equiaxed grains are distributed around these grain bands.

[0120] The inventors found that by adjusting the Cu / Mg ratio to control the stacking fault energy of the Al-Cu-Mg alloy, and adjusting the subsequent annealing process, a combination of Goss+P texture can be formed. When the two texture components have a large torsion or deflection angle with the surrounding grains, the crack usually deflects significantly, and the {111} planes of the two components are close to the direction of the maximum critical shear stress, which relieves the stress concentration during crack propagation. Therefore, the crack propagation resistance is increased, and the fatigue performance is improved. However, the prior inventors disclosed CN103045976A, CN103526140A, CN10358997A, CN108103373B and CN108504915B have excellent fatigue performance, but it is difficult to achieve a material ultimate tensile strength exceeding 480 MPa. Generally, high strength and toughness and excellent fatigue damage resistance are difficult to be compatible, which depends on the strengthening mode in the aluminum alloy, mainly precipitation strengthening and grain boundary strengthening. High density of precipitates and fine grains are beneficial to improve the strength, but these phases hinder the dislocation reciprocating motion during fatigue, thereby reducing the fatigue performance. Furthermore, the above-mentioned disclosed patents detect the fatigue crack propagation rate in the LT direction, and the long diameter direction of the Goss grain is along the RD, which means that most of the grain boundaries are parallel to the RD, that is, the crack propagation direction. Generally, the grain boundary is a weak area, and the grain boundary bonding force is weak, which means that the crack propagation rate in the LT direction is bound to be lower than that in the TL direction. Therefore, how to make the alloy maintain high strength (>495 MPa) and also have a lower fatigue crack propagation rate in the TL direction.

[0121] The present application controls the texture and microstructure of the aged Al-Cu-Mg-Ag alloy by both deformation and micro-alloying, which can make the invented alloy obtain the ultimate tensile strength of more than 500 MPa, while also obtain very excellent anti-fatigue damage performance. The principle of the present application is to regulate the texture, grain structure, coarse second phase particles and reduce the dislocation density of the aged alloy by controlling the deformation temperature and micro-alloying. The specific principle is that under large deformation, deformation at different temperatures can make the alloy form deformed grains with large length-diameter ratio along the RD direction. Under the influence of the deformation temperature, the content of Brass and Copper texture in the deformed matrix can be regulated, and during the subsequent isothermal recrystallization annealing process, since the recrystallized grains grow from the deformed grains, and the grain boundaries restrict the expansion of the recrystallized grains in the TD direction as a barrier, at this time, the grain growth can only continue to expand along the RD, however, there are more growing grains in the RD direction, under the coherent effect of grain boundary migration, finally, Goss or Goss+P grain bands with the grain length-diameter direction along the TD direction are formed. And the surrounding forms a near-equiaxed crystal structure. Therefore, Goss or Goss+P grain bands are formed in the aged alloy, the grain growth direction in the grain band is along the TD direction, and the crack propagation direction is parallel to the RD direction, which means that the grain boundary is perpendicular to the crack propagation direction. In the steady expansion stage, the crack is difficult to penetrate these grains, but expands around the grain boundaries of these grains, which greatly increases the crack propagation process. And to some extent, the formation of surface slip bands is suppressed, the crack initiation caused by slip band immersion is controlled, and the crack propagation in the center of the plate is delayed. The purpose of reducing the crack propagation rate is achieved. Secondly, by controlling the deformation process and micro-alloying, the number density and distribution of coarse Al2CuMg and Al2Cu second phases in the alloy are controlled. Reducing the number density of the two coarse phases promotes the formation of fatigue striation bands in the steady expansion zone, and the dislocation reciprocating sliding becomes easier, thereby inhibiting the generation of micro-cracks between the striation bands. On the basis of controlling the Cu / Mg ratio, Ag element is added, because Ag captures vacancies, the content of ordered dislocations in the aged alloy is finally reduced, further reducing the dislocation sliding resistance and stress concentration acceleration in the fatigue crack propagation process. In addition, under a high Cu / Mg, the Cu atoms in the excess phase after homogenization treatment are maximally dissolved into the matrix, eliminating dendritic segregation, and producing solid solution strengthening effect in the subsequent solid solution treatment. The content of coarse Al2CuMg and Al2Cu phases is reduced, to some extent, the cracking tendency in the plastic deformation process is inhibited, and the continuous hardening ability is produced, so as to contribute to the higher tensile strength. Furthermore, the grain boundaries in the Goss or Goss+P grain band are perpendicular to the tensile stress direction, which further increases the deformation resistance. The plastic damage space, that is, the increase of the yield strength ratio, is due to the inhibition of secondary recrystallization and the formation of atomic clusters by short-time solid solution, thereby obtaining a low yield strength ratio.In addition, by controlling the content of Zn, the consumption of Mg element by Zn is reduced, and the Zn atoms migrate along the grain boundaries during recrystallization, thereby inhibiting the formation of strong Goss texture. With the above principles and the matching technology disclosed in the present application, the Al-Cu-Mg-Ag alloy can have an ultimate tensile strength of more than 500 MPa, a percentage elongation after fracture of 18% to 30%, and very excellent fatigue damage resistance.

[0122] The high-strength aluminum alloy with excellent fatigue damage resistance and the preparation method thereof will be described in detail below with reference to examples and experimental data.

[0123] Example 1

[0124] The chemical composition of the high-strength aluminum alloy in this example is Cu 4.69%, Mg 1.18%, Ag 0.41%, Mn 0.9%, Ti 0.08%, Fe 0.035%, Si 0.035%, Zn 0.07%, and the balance being Al. First, the alloy components are weighed and then melted into an ingot (after melting, the alloy elements may fluctuate, the Cu and Mg elements fluctuate within ±0.2%, the Ag and Mn elements fluctuate within ±0.1%, and the Ti element fluctuates within ±0.02%, and the Fe, Si, and Zn elements are controlled within the designed component range). Then, the ingot is subjected to homogenization treatment at 470°C for 48h and at 505°C for 12h, and then is subjected to hot rolling at a deformation rate of 95.6% at 405°C, with a plate thickness of 2mm. Subsequently, the plate is subjected to isothermal annealing at 340°C for 5h, and then is subjected to solid solution treatment at 485°C for 30min, and then is immediately quenched in water, with a quenching transfer time of 5s, and is naturally aged for 240h.

[0125] The hot-rolled alloy is controlled by Copper texture, with a Brass texture intensity of 6.82, a Copper texture intensity of 8.20, and a S texture intensity of 6.82. The naturally aged alloy is controlled by Goss+P texture, with a Goss texture intensity of 8.01 and a P texture intensity of 8.01. Figure 1 and 2 wherein φ1, Φ, φ2 represent the Euler angles in the Euler space. The naturally aged alloy has a low dislocation content. The yield strength of the naturally aged alloy is 299.9 MPa, the tensile strength is 496.3 MPa, the percentage elongation after fracture is 22.5%, and the yield strength ratio is 0.604. The fatigue crack propagation rate in the TL direction is shown in Table 1, and the instability ΔK is 30.2 MPa*m 1 / 2 .

[0126] wherein φ1, Φ, φ2 represent the Euler angles in the Euler space. The naturally aged alloy has a low dislocation content. The yield strength of the naturally aged alloy is 299.9 MPa, the tensile strength is 496.3 MPa, the percentage elongation after fracture is 22.5%, and the yield strength ratio is 0.604. The fatigue crack propagation rate in the TL direction is shown in Table 1, and the instability ΔK is 30.2 MPa*m 1 / 2 . Figure 3The middle ordinate da / dN (mm / cycle) represents the crack propagation rate, mm / cycle represents the distance of crack propagation per cycle, and the abscissa Delta K (MPa*m1 / 2) represents the stress intensity factor. The room temperature tensile test was performed according to GB 228-87 “Metallic Tensile Test Method” on an MTS Landmark universal testing machine at a rate of 2 mm / min. The yield strength was obtained by clamping the extensometer on the sample, the tensile strength was directly taken from the parameters displayed by the equipment, and the elongation after fracture was measured by a vernier caliper. The fatigue crack propagation rate was performed according to GB / T63982000 on an MTS 810 fatigue testing machine. Before testing, the sample was polished with 180#, 600# and 1000# sandpaper to ensure that the sample surface was free of defects, and the polishing process ensured the uniformity of the sample thickness. The test sample was a standard compact tension (C(T)) sample with a thickness of 2 ± 0.2 mm, and the sample was taken along the T-L direction of the rolled plate. The test methods / detection standards / detection equipment of the yield strength, tensile strength, elongation after fracture and TL direction fatigue crack propagation rate in the subsequent examples and comparative examples were consistent with those in Example 1.

[0127] Table 1 Fatigue crack propagation rate of Example 1 at different ΔK (x10 -4 mm / cycle)

[0128]

[0129] Example 2

[0130] The chemical composition of the high-strength aluminum alloy in this example is Cu 4.87%, Mg 1.27%, Ag 0.53%, Mn 0.71%, Ti 0.07%, Fe 0.045%, Si 0.032%, Zn 0.06%, and the balance is Al. First, the alloy was charged according to the above alloy composition and then melted into an ingot. After melting, the alloy elements will fluctuate, the Cu and Mg elements will fluctuate within ±0.2%, the Ag and Mn elements will fluctuate within ±0.1%, and the Ti element will fluctuate within ±0.02%, and the Fe, Si and Zn elements will be controlled within the designed composition range. Subsequently, the ingot was subjected to homogenization treatment at 475°C / 36h + 500°C / 24h, and then hot rolling at 490°C with a deformation of 90%, and the plate thickness was 3mm. Subsequently, isothermal annealing at 390°C for 4h, followed by solid solution treatment at 495°C for 15min, followed by immediate water quenching, quenching transfer time 3s, natural aging 260h.

[0131] The hot-rolled alloy is controlled by Brass texture, the Brass texture strength of the hot-rolled alloy is 8.68, the Copper texture strength is 4.32, and the S texture strength is 4.32. The natural aging alloy is controlled by Goss + Cube texture, the Goss texture strength is 9.28, and the Cube texture strength is 9.28.Figure 5 and 6 PHI, PHI, Constant PHI2 in FIG. 1, FIG. 2 and FIG. 3 represent φ1, Φ, φ2 angles in Euler space, there are local ordered spiral dislocations in natural aging alloy, and the density is low. The yield strength of the natural aging alloy is 311.6 MPa, the tensile strength is 507.4 MPa, the elongation after fracture is 23%, and the yield strength ratio is 0.61. The fatigue crack propagation rate of the TL direction is shown in Table 2, and the instability ΔK = 30.4 MPa*m 1 / 2 . The Figure 7 The vertical coordinate da / dN (mm / cycle) in FIG. 1, FIG. 2 and FIG. 3 represents the crack propagation rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0132] Table 2 Fatigue crack propagation rate of Example 2 alloy under different ΔK (x10 -4 mm / cycle)

[0133]

[0134] Example 3

[0135] The chemical composition of the high-strength aluminum alloy of this example is Cu 5.32%, Mg 1.38%, Ag 0.89%, Mn 0.62%, Ti 0.06%, Fe 0.05%, Si 0.035%, Zn 0.076%, and the balance is Al. First, the alloy is prepared according to the above alloy composition, and then melted into an ingot. After melting, the alloy elements will fluctuate, the Cu and Mg elements fluctuate within ±0.2%, the Ag and Mn elements fluctuate within ±0.1%, and the Ti element fluctuates within ±0.02%. The Fe, Si and Zn elements are controlled within the designed composition range. Subsequently, the ingot is subjected to homogenization treatment at 465°C / 36h+495°C / 48h, and then hot rolling at 435°C with a deformation of 93%, and the hot rolling plate thickness is 2mm. Subsequently, isothermal annealing at 360°C for 5h, and then solid solution treatment at 503°C for 3min, followed by immediate water quenching, quenching transfer time 5s, and natural aging for 280h.

[0136] The yield strength of the natural aging alloy is 319.8 MPa, the tensile strength is 512.7 MPa, the elongation after fracture is 23.2%, and the yield strength ratio is 0.637. The fatigue crack propagation rate of the TL direction is shown in Table 3, and the instability ΔK = 30.84 MPa*m 1 / 2 . The Figure 9 The vertical coordinate da / dN (mm / cycle) in FIG. 1, FIG. 2 and FIG. 3 represents the crack propagation rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0137] Table 3 Fatigue crack propagation rate of Example 3 alloy under different ΔK (x10-4 mm / cycle)

[0138]

[0139] Example 4

[0140] The chemical composition of the high-strength aluminum alloy of this example is Cu 5.05%, Mg 1.3%, Ag 0.65%, Mn 0.82%, Ti 0.06%, Fe 0.05%, Si 0.04%, Zn 0.07%, and the balance being Al. The alloy ingot was first prepared by melting and casting according to the above alloy composition. After melting, the alloy elements fluctuate within the following ranges: Cu and Mg fluctuate within ±0.2%, Ag and Mn fluctuate within ±0.1%, and Ti fluctuates within ±0.02%, while Fe, Si, and Zn are controlled within the designed composition range. The ingot was then subjected to homogenization treatment at 480°C for 36h and at 495°C for 24h, followed by hot rolling at 455°C with a deformation of 85%, and then directly cold rolling at 45% with a cold-rolled sheet thickness of 1.8mm. Subsequently, the sheet was subjected to isothermal annealing at 365°C for 6h, followed by solution treatment at 500°C for 10min, and then immediately water quenching with a quenching transfer time of 7s, and natural aging for 280h.

[0141] The cold-rolled alloy is controlled by Brass texture with a Brass texture intensity of 9.64, a Copper texture intensity of 3.18, and a S texture intensity of 3.18. The naturally aged alloy is controlled by Goss+Cube texture with a Goss texture intensity of 7.1 and a Cube texture intensity of 4.1. Figure 10 and 11 wherein φ1, Φ, φ2 represent the Euler angles in the Euler space. In the naturally aged alloy, only a small amount of ordered dislocation loops and kinked dislocations exist. The microstructure contains Goss+Cube grain bands, the grains inside the grain bands have a long diameter parallel to the TD direction, the grain boundaries are perpendicular to the RD direction, and equiaxed grains are distributed around the grain bands. The yield strength of the naturally aged alloy is 326.1MPa, the tensile strength is 512.1MPa, the elongation after fracture is 21.1%, and the yield strength ratio is 0.637. The fatigue crack propagation rate of the alloy in the TL direction is shown in Table 4, and the instability ΔK = 37.7MPa*m 1 / 2 Figure 12 wherein the ordinate da / dN (mm / cycle) represents the crack propagation rate, and the abscissa DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0142] Table 4 Fatigue crack propagation rate of the alloy of Example 4 at different ΔK (×10 -4 mm / cycle)

[0143] ​

[0144] Example 5

[0145] The chemical composition of the high strength aluminum alloy of this example is Cu 5.10%, Mg 1.30%, Ag 0.45%, Mn 0.65%, Ti 0.09%, Fe 0.045%, Si 0.05%, Zn 0.05%, balance Al. The alloy was first charged according to the above alloy composition and then melted into ingot. After melting, the alloy elements will fluctuate, the Cu and Mg elements fluctuate within ±0.2%, the Ag and Mn elements fluctuate within ±0.1%, and the Ti element fluctuates within ±0.02%, and the Fe, Si and Zn elements are controlled within the designed composition range. The ingot was then subjected to homogenization treatment at 485°C / 24h + 495°C / 36h, and then hot rolling at 465°C with a deformation of 91%, and the hot rolled plate was annealed at 300°C / 3h, and then cold rolling at 30%, and the cold rolled plate thickness was 2mm. Subsequently, isothermal annealing at 370°C for 4h, followed by solution treatment at 505°C for 6min, followed by immediate water quenching, quenching transfer time 10s, natural aging 265h.

[0146] The natural aging alloy is controlled by Goss + Cube texture, the Goss texture strength is 10.7, and the Cube texture strength is 8.2. The Figure 15 PHI, Constant PHI2 in the formula represent the φ1, Φ, φ2 angles in the Euler space. Compared with Example 3, the Goss and Cube texture strengths of the aged alloy are significantly increased. There are low-density dislocation lines and dislocation tangles in the natural aging alloy. The yield strength of the natural aging alloy is 321.6MPa, the tensile strength is 505.1MPa, the elongation after fracture is 18.9%, and the yield strength ratio is 0.637. The fatigue crack propagation rate in the TL direction is shown in Table 5. The instability ΔK = 39.4MPa*m 1 / 2 The Figure 16 The ordinate da / dN (mm / cycle) in the formula represents the crack propagation rate, and the abscissa DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0147] Table 5 Fatigue crack propagation rate of Example 5 alloy at different ΔK (x10 -4 mm / cycle)

[0148]

[0149] Example 6

[0150] The chemical composition of the high-strength aluminum alloy of this example is Cu 4.93%, Mg 1.24%, Ag 0.7%, Mn 0.59%, Ti 0.075%, Fe 0.038%, Si 0.041%, Zn 0.062%, and the balance being Al. First, the alloy is batched according to the above alloy composition and then melted into an ingot. After melting, the alloy elements will fluctuate, with the Cu and Mg elements fluctuating within ±0.2%, the Ag and Mn elements fluctuating within ±0.1%, and the Ti element fluctuating within ±0.02%, and the Fe, Si, and Zn elements being controlled within the designed composition range. Subsequently, the ingot is subjected to homogenization treatment at 490°C / 36h + 490°C / 36h, and then is subjected to hot rolling at 480°C with a deformation of 91.5% and a plate thickness of 2.5mm. Subsequently, the plate is subjected to isothermal annealing at 415°C for 3h, and then is subjected to solid solution treatment at 495°C for 8min, and then is immediately water quenched, with a quenching transfer time of 11s, an artificial aging temperature of 155°C, and an aging time of 3h.

[0151] In the artificial aging state, the dislocations in the alloy almost disappear, and there are low-density atomic clusters. The yield strength of the artificial aging alloy is 316.3MPa, the tensile strength is 506.5MPa, the elongation after fracture is 23.8%, and the yield strength ratio is 0.624. The fatigue crack propagation rate of the alloy in the TL direction is shown in Table 6, and the instability ΔK = 35.1MPa*m 1 / 2 . The Figure 18 The ordinate da / dN (mm / cycle) represents the crack propagation rate, and the abscissa DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0152] Table 6 Fatigue crack propagation rate of the alloy of Example 6 at different ΔK (x10 -4 mm / cycle)

[0153]

[0154] Example 7

[0155] The chemical composition of the high-strength aluminum alloy of this example is Cu 4.78%, Mg 1.4%, Ag 0.6%, Mn 0.77%, Ti 0.083%, Fe 0.032%, Si 0.034%, Zn 0.043%, and the balance Al. First, the alloy is charged according to the above alloy composition and then melted into an ingot. After melting, the alloy elements will fluctuate, the Cu and Mg elements will fluctuate within ±0.2%, the Ag and Mn elements will fluctuate within ±0.1%, and the Ti element will fluctuate within ±0.02%. The Fe, Si, and Zn elements are controlled within the designed composition range. Subsequently, the ingot is subjected to homogenization treatment at 485°C / 24h + 490°C / 48h, and then hot-rolled at 470°C with a deformation of 87% and a plate thickness of 3.2mm. Subsequently, the plate is subjected to isothermal annealing at 400°C for 4h, and then solid solution treatment at 490°C for 20min, followed by water quenching immediately. The quenching transfer time is 6s, the artificial aging temperature is 173°C, and the aging time is 1h.

[0156] In the artificial aging state, the dislocations in the alloy almost disappear, and there are low-density atomic clusters. Compared with Example 5, the size of the atomic clusters slightly increases. The yield strength of the artificial aging alloy is 341.4MPa, the tensile strength is 518.6MPa, the elongation after fracture is 23.4%, and the yield strength ratio is 0.658. The fatigue crack propagation rate of the alloy in the TL direction is shown in Table 7, and the instability ΔK = 31.2MPa*m 1 / 2 . Attach Figure 20 The middle ordinate da / dN (mm / cycle) represents the crack propagation rate, and the abscissa DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0157] Table 7 Fatigue crack propagation rate of the alloy of Example 7 at different ΔK (×10 -4 mm / cycle)

[0158]

[0159] Comparative Example 1

[0160] Comparative Example 1 has the following chemical composition: Cu 4.1%, Mg 1.7%, Mn 0.71%, Ti 0.07%, Fe 0.045%, Si 0.032%, Zn 0.06%, with the balance being Al. First, the alloying elements were melted and cast into ingots according to the above composition. After melting, the alloying elements fluctuated: Cu and Mg fluctuated within ±0.2%, Mn within ±0.1%, and Ti within ±0.02%. Fe, Si, and Zn were controlled within the designed composition range. The ingots were then homogenized at 475℃ for 36h followed by 500℃ for 24h, and then hot-rolled at 490℃ with a deformation of 90%, resulting in a plate thickness of 3mm. Subsequently, it was isothermally annealed at 390℃ for 4h, followed by solution treatment at 495℃ for 15min, then immediately water quenched with a quenching transfer time of 3s, and naturally aged for 260h.

[0161] The hot-rolled sheet has a Brass texture strength of 12.23, a Copper texture strength of 1.99, and an S texture strength of 1.99. The naturally aged alloy is controlled by a Goss+Cube texture, with both the Goss and Cube texture strengths being 7.87. (See attached image.) Figure 22 and 23 In the figure, Phi1, Pi, and Constant Phi2 represent angles φ1, φ2 in Euler space. The naturally aged alloy consists of highly ordered dislocation loops with a high dislocation density. The naturally aged alloy has a yield strength of 305.7 MPa, a tensile strength of 477.5 MPa, an elongation after fracture of 23.6%, and a yield-to-tensile ratio of 0.64. It can be seen that although the naturally aged alloy in Comparative Example 1 has good fatigue performance, its ultimate tensile strength is far below 500 MPa. The fatigue crack propagation rate in the TL direction is shown in Table 8, with an instability ΔK = 33.5 MPa*m. 1 / 2 Appendix Figure 24 The vertical axis da / dN (mm / cyc le) represents the crack propagation rate, and the horizontal axis De ltaK (MPa*m1 / 2) represents the stress intensity factor.

[0162] Table 8. Fatigue crack propagation rate of alloy in Comparative Example 1 at different ΔK (×10) -4 mm / cycle)

[0163]

[0164]

[0165] Comparative Example 2

[0166] The chemical composition of the alloy in Comparative Example 2 is Cu 4.83%, Mg 1.25%, Mn 0.6%, Ti 0.08%, Fe 0.041%, Si 0.036%, Zn 0.08%, with the balance being Al. This alloy has a high Cu / Mg ratio but no Ag is added. The alloy was first prepared according to the above composition and then melted into an ingot. After melting, the alloy elements fluctuated: Cu and Mg fluctuated within ±0.2%, Mn within ±0.1%, and Ti within ±0.02%. Fe, Si, and Zn were controlled within the designed composition range. The ingot was then homogenized at 475℃ / 34h + 500℃ / 20h, and then hot-rolled at 486℃ with a deformation of 91%, resulting in a plate thickness of 2.8mm. It was then isothermal annealed at 385℃ for 4 hours, followed by solution treatment at 495℃ for 15 minutes, and then immediately water quenched with a quenching transfer time of 4 seconds and natural aging for 260 hours.

[0167] The hot-rolled sheet has a Brass texture strength of 11.31, a Copper texture strength of 5.63, and an S texture strength of 3.73. The naturally aged alloy is controlled by a Goss+Cube texture, with both the Goss and Cube texture strengths being 9.96. (See attached image.) Figure 26 and 27 In the equation, Phi1, Pi, and Constant Phi2 represent angles φ1, φ2 in Euler space. The naturally aged alloy consists of highly ordered dislocation loops with a high dislocation density. The naturally aged alloy exhibits a yield strength of 311.9 MPa, a tensile strength of 494.8 MPa, an elongation after fracture of 21.9%, and a yield-to-tensile ratio of 0.63. It can be seen that the naturally aged alloy in Comparative Example 2 has poor fatigue performance, especially under large ΔK conditions, where the crack propagation rate is significantly accelerated, and the tensile strength has not yet exceeded 500 MPa. ΔK = 30 MPa * m 1 / 2 The crack propagation rate is as high as 3.5 × 10⁻⁶. -3 mm / cycle. The fatigue crack propagation rate in the TL direction is shown in Table 9, with an instability ΔK = 30.1 MPa*m. 1 / 2 Appendix Figure 28 The vertical axis da / dN (mm / cyc le) represents the crack propagation rate, and the horizontal axis De ltaK (MPa*m1 / 2) represents the stress intensity factor.

[0168] Table 9. Fatigue crack propagation rates of alloy 2 under different ΔK (×10) -4 mm / cycle)

[0169]

[0170] Comparative Example 3

[0171] The chemical composition of the high-strength aluminum alloy of this embodiment is Cu 4.90%, Mg 1.25%, Mn 0.60%, Ti 0.072%, Fe 0.038%, Si 0.045%, Zn 0.06%, and the balance being Al. First, the above alloy composition is batched and then smelted into an ingot. After smelting, the alloy elements will fluctuate, the Cu and Mg elements will fluctuate within ±0.2%, the Mn element will fluctuate within ±0.1%

[0172] ±0.02%, and the Fe, Si, and Zn elements are controlled within the designed composition range. Subsequently, the ingot is subjected to homogenization treatment at 490°C / 35h+490°C / 36h, and then is subjected to hot rolling at 480°C with a deformation of 91% and a plate thickness of 2.4mm. Subsequently, the plate is subjected to isothermal annealing at 410°C for 3h, and then is subjected to solid solution treatment at 495°C for 7min, and then is immediately water quenched, with a quenching transfer time of 10s, an artificial aging temperature of 159°C, and an aging time of 3h.

[0173] The artificial aging alloy is also composed of highly ordered dislocation loops, and has a high dislocation density. The yield strength of the artificial aging alloy is only 280.2MPa, the tensile strength is only 469.4MPa, the elongation after fracture is 23.7%, and the yield strength ratio is 0.597. After artificial aging, the strength of the Comparative Example 3 alloy decreases sharply, and the ultimate tensile strength is far from 500MPa. Although the crack propagation rate is slow at a low ΔK, the alloy has already been unstable and fractured at ΔK=26.6MPa*m 1 / 2 The fatigue crack propagation rate of the alloy of Comparative Example 3 at different ΔK is shown in Table 10. The Figure 30 vertical ordinate da / dN (mm / cycle) represents the crack propagation rate, and the horizontal ordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0174] Table 10 Fatigue crack propagation rate (×10 -4 mm / cycle) of Comparative Example 3 alloy at different ΔK

[0175]

[0176] Comparative Example 4

[0177] The chemical composition of the alloy of Comparative Example 4 is Cu 4.95%, Mg 1.69%, Ag 0.38%, Mn 0.6%, Ti 0.08%, Fe 0.045%, Si 0.045%, Zn 0.09%, and the balance Al. The alloy ingot is prepared by first charging the alloying elements according to the above composition and then melting. After melting, the alloying elements fluctuate within the following ranges: Cu and Mg fluctuate within ±0.2%, Ag and Mn fluctuate within ±0.1%, and Ti fluctuates within ±0.02%, and Fe, Si and Zn are controlled within the designed composition range. The ingot is then subjected to homogenization treatment at 490°C / 36h + 495°C / 24h, and then hot-rolled at 480°C with a deformation of 90% to a sheet thickness of 2mm. The sheet is then subjected to isothermal annealing at 375°C for 4h, and then solid solution treatment at 493°C for 11min, followed by water quenching. The quenching transfer time is 6s, and the natural aging time is 255h.

[0178] The hot-rolled alloy is controlled by Brass texture, and the Brass texture intensity of the hot-rolled alloy is 10.71, the Copper texture intensity is 1.74, and the S texture intensity is 1.74. The naturally aged alloy is controlled by Goss + Cube texture, and the Goss texture intensity is 5.15, and the Cube texture intensity is 5.15. Figure 32 and 33 wherein φ1, Φ, φ2 represent the Euler space angles of φ1, Φ, φ2. The yield strength of the naturally aged alloy is 303.9MPa, the tensile strength is 482.3MPa, the elongation after fracture is 18.0%, and the yield strength ratio is 0.630. The naturally aged alloy is composed of high-density dislocation lines. The fatigue damage index in the TL direction is shown in Table 11. The instability ΔK = 28.5MPa*m 1 / 2 Compared with Example 2, the Cu / Mg is reduced, and the Mg content is increased, which results in a decrease in the Goss texture intensity of the aged alloy, and a sharp increase in the content of coarse second phase in the matrix. Not only is the ultimate tensile strength far from 500MPa, but also the fatigue performance is very poor. The fatigue crack propagation rate in the TL direction is shown in Table 11. Figure 34 The ordinate da / dN (mm / cycle) in Table 11 represents the crack propagation rate, and the abscissa DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0179] Table 11 Fatigue crack propagation rate (×10 -4 mm / cycle) of Comparative Example 4 at different ΔK

[0180]

[0181] The fatigue crack propagation rate of Comparative Example 4 is shown in Table 11. Figures 1-4It can be seen that after homogenization at 470℃ / 48h + 505℃ / 12h + 95.6% reduction at 405℃, the hot-rolled texture is controlled by high-strength Copper, S, and Brass textures, with strengths of 8.20, 6.82, and 6.82 respectively, among which the Copper texture has the highest strength. After annealing at 340℃ / 5h + solution treatment at 485℃ / 30min + water quenching (transfer time 5s) + at least 240h of natural aging, the hot-rolled plate is controlled by high-strength Goss and P textures, with strengths of 8.01 for both Goss and P textures and a combined strength of 16.02. The steady-state expansion rate of the alloy in Example 1 is relatively low, at ΔK = 30MPa*m. 1 / 2 At that time, the expansion rate was only 15.2 × 10 -4 mm / cycle. The naturally aged matrix has a low content of internal dislocations, with only a small number of coiled dislocations.

[0182] Combined with appendix Figures 5-8 It can be seen that after homogenization at 475℃ / 36h + 500℃ / 24h + 90% reduction at 490℃, the hot-rolled texture is controlled by a high-strength Brass texture with a strength of 8.68, the Copper texture strength is 4.32, and the S texture strength is 4.32. After annealing at 390℃ / 4h + solution treatment at 495℃ / 15min + water quenching (transfer time 3s) + at least 260h of natural aging, the hot-rolled plate is controlled by high-strength Goss and Cube textures, with strengths of both Goss and Cube textures at 9.28 and a combined strength of 18.56. Example 2 alloy at ΔK = 30MPa*m 1 / 2 At that time, the expansion rate was 18.4 × 10 -4 mm / cycle. The naturally aged matrix consists of long coiled dislocations and dislocation loops.

[0183] from Figure 9 As can be seen from the process of homogenization at 465℃ / 36h + 495℃ / 48h + 93% reduction at 435℃ + annealing at 360℃ / 5h + solution treatment at 503℃ / 3min + water quenching (transfer time 3s) + natural aging for at least 260h, the alloy in Example 3 was subsequently subjected to ΔK = 30MPa*m. 1 / 2 At that time, the expansion rate was 26.9 × 10 -4 mm / cycle.

[0184] Combined with appendix Figures 10-14It can be seen that after homogenization at 480°C / 36h + 495°C / 24h + hot rolling at 455°C with 85% reduction + 45% cold rolling the texture is controlled by a high intensity Brass texture with Brass texture intensity of 9.64, Copper texture intensity of 3.18 and S texture intensity of 3.18. After annealing at 365°C / 6h + solutionizing at 500°C / 10min + water quenching (transfer time 7s) + at least 280h natural aging the cold rolled sheet is controlled by a high intensity Goss texture with Goss texture intensity of 7.1, Cube texture intensity of 4.1 and combined intensity of 11.2. The alloy of example 3 has a crack propagation rate of 19.2 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 mm / cycle at ΔK = 30 MPa*m

[0185] In combination with the attached figures Figures 15-17 It can be seen that after homogenization at 485°C / 36h + 495°C / 36h + hot rolling at 465°C with 91% reduction + annealing at 300°C / 3h + 30% cold rolling + annealing at 370°C / 4h + solutionizing at 505°C / 6min + water quenching (transfer time 10s) + at least 265h natural aging the texture is controlled by a high intensity Goss + Cube texture with Goss texture intensity of 10.7, Cube texture intensity of 8.2 and combined intensity of 18.9. The alloy of example 4 has a crack propagation rate of 22.7 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 mm / cycle at ΔK = 30 MPa*m

[0186] In combination with the attached figures Figures 18-19 It can be seen that after homogenization at 490°C / 36h + 490°C / 36h + hot rolling at 480°C with 91.5% reduction + annealing at 415°C / 3h + solutionizing at 495°C / 8min + water quenching (transfer time 11s) + artificial aging at 155°C / 3h the alloy of example 5 has a crack propagation rate of 15.8 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 mm / cycle at ΔK = 30 MPa*m

[0187] In combination with the attached figures Figures 20-21It can be seen that the Example 6 alloy after homogenization at 485°C / 24h + 490°C / 48h + hot rolling at 470°C with 87% reduction + annealing at 400°C / 4h + solutionizing at 495°C / 20min + water quench (transfer time 6s) + artificial aging at 173°C / 1h has an expansion rate of 29.1 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 The matrix in the artificially aged state consists of large size, high density Cu-Mg atom clusters with dislocations essentially eliminated.

[0188] In conjunction with the attached Figures 22-25 It can be seen that the Comparative Example 1 alloy after homogenization at 475°C / 36h + 500°C / 24h + hot rolling at 490°C with 90% reduction + annealing at 390°C / 4h + solutionizing at 495°C / 15min + water quench (transfer time 3s) + natural aging for at least 260h has an expansion rate of 14.9 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 The matrix in the naturally aged state consists of high density dislocation loops.

[0189] In conjunction with the attached Figures 26-29 It can be seen that the Comparative Example 2 alloy after homogenization at 475°C / 34h + 500°C / 20h + hot rolling at 486°C with % reduction + annealing at 385°C / 4h + solutionizing at 495°C / 15min + water quench (transfer time 4s) + natural aging for at least 260h has an expansion rate of 35 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 The matrix in the naturally aged state consists of high density dislocation loops.

[0190] In conjunction with the attached Figures 30-31 It can be seen that the Comparative Example 3 alloy after homogenization at 490°C / 35h + 490°C / 36h + hot rolling at 480°C with 91% reduction + annealing at 410°C / 3h + solutionizing at 495°C / 7min + water quench (transfer time 10s) + artificial aging at 159°C / 3h has an expansion rate of 22.5 x 10 1 / 2The time has been broken. The artificial aging state matrix is composed of high density dislocation ring.

[0191] In combination with the Figures 32-35 It can be seen that the hot-rolled texture of the 90% reduction at 480°C is controlled by high strength Brass texture with a strength of 10.71, Copper texture strength of 1.74, and S texture strength of 1.74. After the hot-rolled plate is annealed at 375°C for 4h, solid-solved at 493°C for 11 min, water quenched (transferring time of 6s), and naturally aged for at least 255h, it is controlled by Goss + Cube texture, with Goss texture strength of only 5.15, Cube texture strength of only 5.15, and combined strength of 10.30. The expansion rate of the steady-state expansion zone of the alloy of Comparative Example 4 is relatively high, and the crack expansion rate is as high as 71.6 x 10 1 / 2 mm / cycle at ΔK = 25 MPa*m - 4 The matrix of the natural aging state is composed of high density dislocation lines.

[0192] In combination with the Figures 1-21 It can be seen that the tensile strength of the alloy of the present application can be higher than 495 MPa after the process treatment of homogenization, hot-rolling or hot-rolling and cold-rolling, annealing, solid-solution and subsequent aging of the alloy in the scope of claims, and the fatigue crack expansion rate of the alloy is lower than 30 x 10 1 / 2 mm / cycle at ΔK = 30 MPa*m -4 mm / cycle at ΔK = 30 MPa*m. Excellent properties of high strength and fatigue resistance are presented.

[0193] In combination with the Figures 5-8 It can be seen from Examples 22-29 that when the Cu content in the alloy is lower than 4.5%, the Mg content is higher than 1.4%, and no Ag element is added, the Cu / Mg ratio is lower than 3≤Cu / Mg≤4, although the Brass texture strength of the deformed state is relatively high, the Goss and Cube texture strengths are decreased. And the dislocation density of the natural aging state is significantly increased. Although the fatigue performance is improved, the tensile strength is far from 495 MPa. In addition, keeping the Cu and Mg element contents close and not adding Ag element leads to the formation of high density dislocation ring in the natural aging state alloy, and there is no significant difference in the Goss and Cube texture strengths of the natural aging state, but the tensile strength of the alloy is also lower than 495 MPa, and the fatigue crack expansion rate at high ΔK is significantly increased.

[0194] In combination with the Figures 10-17 It can be seen that, before cold rolling, the hot-rolled plate is annealed at 300°C for 3h, compared with the plate without annealing before cold rolling, the Brass texture is enhanced, and the Goss texture strength of the natural aging state is significantly enhanced, thereby improving the instability ΔK.

[0195] Fig. 1 shows the microstructure of the alloy of Example 1. Figures 18-19 As can be seen from Figs. 2-5, the Cu and Mg element contents are kept close to each other, and no Ag element is added. The alloy without Ag element after underaging still consists of high density dislocation loops, and the tensile strength is far lower than 495 MPa, and the fatigue performance is also significantly reduced. The alloy without Ag element has already been unstable and fractured at ΔK = 26.6 MPa*m 1 / 2 .

[0196] Fig. 6 shows the microstructure of the alloy of Example 7. Figures 5-8 As can be seen from Figs. 6-9, when the Cu content in the alloy is not much different, and the Mg content is higher than 1.4%, the high density dislocation of disorder is formed in the naturally aged alloy, and the Goss and Cube texture strength is significantly reduced. In turn, not only the tensile strength of the naturally aged alloy is significantly reduced, but also the fatigue performance is greatly deteriorated. The crack propagation rate is as high as 71.6 x 10 1 / 2 mm / cycle at ΔK = 25 MPa*m -4 . However, when the Cu and Mg elements are within the scope of the claims of the present application, the crack propagation rate is only 8.9 x 10 1 / 2 mm / cycle at ΔK = 25 MPa*m -4 .

[0197] As can be seen from the tensile mechanical properties and fatigue crack propagation rates of Comparative Examples 1-4 and Examples 1-7, the alloy of Examples 1-7 still has very excellent fatigue performance under the premise of ensuring high strength (> 495 MPa).

[0198] Finally, it should be noted that the terms "comprising", "including", or any other variation thereof, are intended to cover a non-exclusive inclusion, such that processes, methods, articles, or apparatuses that comprise a list of elements are not limited to those elements, but can include other elements not expressly listed or inherent to such processes, methods, articles, or apparatuses.

[0199] Although preferred embodiments of the present application have been described, those skilled in the art can make further changes and modifications to these embodiments once they have the basic inventive concept. Therefore, the appended claims are intended to be interpreted as including the preferred embodiments and all changes and modifications falling within the scope of the present application.

[0200] Obviously, those skilled in the art can make various modifications and variations to the present application without departing from the spirit and scope of the present application. Thus, if these modifications and variations of the present application fall within the scope of the claims of the present application and their equivalents, the present application is also intended to include these modifications and variations.

Claims

1. A high-strength aluminum alloy with excellent resistance to fatigue damage, characterized in that, The chemical composition of the high-strength aluminum alloy, by mass fraction, is as follows: Cu: 4.5–5.5%, Mg: 1.0–1.4%, Ag: 0.2–1%, Mn: 0.5–1%, Ti: 0.06–0.10%, Fe<0.06%, Si<0.06%, Zn<0.1%; the balance is Al, wherein the mass ratio of Cu to Mg is (4.78:1.4)–4; The microstructure of the high-strength aluminum alloy contains Goss+Cube or Goss+P grain bands, the major axis of the grains inside the grain bands is parallel to the TD direction, the grain boundaries are perpendicular to the RD direction, and equiaxed grains are distributed around the grain bands. The high-strength aluminum alloy has a yield strength of 300–350 MPa, an ultimate tensile strength of 495–520 MPa, and an elongation after fracture of 18–30%. The fatigue damage index of the high-strength aluminum alloy in the TL direction is: ΔK = 20 MPa * m 1 / 2 da / dN < 5 × 10 -4 mm / cycle; ΔK=30MPa*m 1 / 2 da / dN < 3 × 10 -3 mm / cycle; ΔK=33MPa*m 1 / 2 da / dN < 4 × 10 -3 mm / cycle, instability ΔK can reach 39MPa*m 1 / 2 above.

2. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance as described in claim 1, characterized in that, The preparation method includes: Ingots are obtained through smelting and casting processes; The ingot is subjected to a first-stage homogenization treatment, followed by a second-stage homogenization treatment at a higher temperature to obtain a homogenized ingot. The homogenized ingot is subjected to large deformation hot rolling to obtain a hot-rolled plate; Alternatively, the homogenized ingot can be subjected to large deformation hot rolling, followed by annealing and then cold rolling or direct cold rolling to obtain a cold-rolled sheet. The hot-rolled plate or the cold-rolled plate is isothermally annealed to obtain an annealed aluminum alloy plate; The annealed aluminum alloy plate is subjected to high-temperature short-time solution treatment, water quenching and aging treatment to obtain a high-strength aluminum alloy. The chemical composition of the ingot, by mass fraction, is as follows: Cu: 4.5–5.5%, Mg: 1.0–1.4%, Ag: 0.2–1%, Mn: 0.5–1%, Ti: 0.06–0.10%, Fe<0.06%, Si<0.06%, Zn<0.1%; balance Al, 3≤Cu / Mg≤4.

3. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 2, characterized in that, The process of performing a primary homogenization treatment on the ingot, followed by a secondary homogenization treatment at a higher temperature, to obtain a homogenized ingot, specifically includes: The ingot is subjected to a first-stage homogenization treatment at 465–490°C for 24–48 hours. After the first homogenization treatment, the temperature is raised to 490℃~505℃ for the second homogenization treatment, and the homogenization time is 12~48h to obtain homogenized ingots.

4. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 2, characterized in that, The step of hot rolling the homogenized ingot with large deformation to obtain a hot-rolled plate specifically includes: The homogenized ingot is subjected to large deformation hot rolling at 390-490°C, with a total rolling deformation of more than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm.

5. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 2, characterized in that, The process of hot rolling the homogenized ingot with large deformation, followed by annealing and then cold rolling or direct cold rolling to obtain a cold-rolled sheet specifically includes: The homogenized ingot is subjected to large deformation hot rolling at 390-490℃, with a total rolling deformation of more than 85%, to obtain a hot-rolled plate with a thickness of 2-6mm. The hot-rolled plate is then directly subjected to cold rolling, with a cold rolling deformation of ≤50%, to obtain a cold-rolled plate. Alternatively, the homogenized ingot is subjected to large deformation hot rolling at 390–490°C, with a total rolling deformation of more than 85%, to obtain a hot-rolled plate with a thickness of 2–6 mm. The hot-rolled plate is then annealed at a temperature of 260°C–330°C for 2–4 hours. Subsequently, the annealed plate is cold-rolled with a cold rolling deformation of ≤50% to obtain a cold-rolled plate.

6. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 2, characterized in that, The step of isothermal annealing the hot-rolled plate or the cold-rolled plate to obtain an annealed aluminum alloy plate specifically includes: The hot-rolled plate or the cold-rolled plate is isothermally annealed to obtain an annealed aluminum alloy plate; Specifically, when the hot rolling temperature of large deformation is below 450℃, the isothermal annealing temperature is controlled at 320℃~370℃ and the annealing time is 2~5h. When the hot rolling temperature of large deformation is above 450℃, the isothermal annealing temperature is controlled at 340℃~420℃ and the annealing time is 3~6h.

7. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 2, characterized in that, The process of subjecting the annealed aluminum alloy plate to high-temperature short-time solution treatment, water quenching, and aging treatment to obtain a high-strength aluminum alloy specifically includes: The annealed aluminum alloy plate is subjected to high-temperature short-time solution treatment at a temperature of 475-505℃ for 3-30 minutes, followed by water quenching with water as the medium and a quenching transfer time of ≤15 seconds. After aging treatment, a high-strength aluminum alloy is obtained. When the aging process adopts the natural aging system, the natural aging time is ≥240h; when the artificial aging system is adopted, the artificial aging temperature is 140℃~180℃ and the aging time is 1~5h.

8. The application of the high-strength aluminum alloy with excellent fatigue damage resistance as described in claim 1 in the manufacture of structural components for aerospace vehicles.

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