High-strength aluminum alloy with excellent fatigue damage resistance and preparation method thereof

Abstract

The invention provides a high-strength aluminum alloy with excellent fatigue damage resistance, and belongs to the technical field of non-ferrous metal material preparation, and the high-strength aluminum alloy comprises the following chemical components in percentage by mass: 4.5-5.5% of Cu, 1.0-1.4% of Mg, 0.2-1% of Ag, 0.5-1% of Mn, 0.06-0.10% of Ti, and the balance of Fe. 0.06%, Silt; 0.06% of Zn < lt >; 0.1%; cu / Mg is greater than or equal to 3 and less than or equal to 4; the microstructure of the high-strength aluminum alloy comprises a Goss + Cube or Goss + P grain belt, the long diameter of grains in the grain belt is parallel to the TD direction, the grain boundary is perpendicular to the RD direction, and equiaxed grains are distributed around the grain belt. The high-strength aluminum alloy has the advantages of high strength and toughness, excellent fatigue damage resistance, reasonable material component design, simple preparation process flow and low industrial production cost. The invention further provides a preparation method of the high-strength aluminum alloy with the excellent fatigue damage resistance.

Description

Technical Field

[0001] The invention belongs to the technical field of nonferrous metal material preparation, and particularly relates to a high-strength aluminum alloy with excellent fatigue damage resistance and a preparation method thereof. Background Art

[0002] The 2000 series aluminum alloy has medium strength and good fatigue resistance, and is the main candidate material for aircraft structural parts and aircraft skins. The typical alloy is the 2x24 series alloy developed by Alcoa in the United States. For example, the 2524 alloy is an improved version based on the 2024 alloy. By controlling the content of Cu and Mg and reducing the content of Fe and Si, the fracture toughness and fatigue performance are effectively improved, but the strength of the alloy is reduced. The fourth-generation aluminum-lithium alloy 2060-T8E30 developed by Alcoa for my country's C919 aircraft has a tensile strength of more than 500MPa and relatively excellent fatigue damage resistance. It is used to replace the traditional 2x24 series alloy. However, the preparation of this alloy is difficult, and the preparation cost is significantly increased compared to the 2x24 series alloy. 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 of materials that affect fatigue performance mainly include: grain size, second phase particles, grain boundary structure and texture. The authorized invention patent technology that improves fatigue performance through texture component design is mainly CN103045976A, CN103526140A, CN10358997A, CN108103373B and CN108504915B announced by the inventors in recent years. The fatigue resistance of Al-Cu-Mg alloys announced by these technologies is relatively excellent, but the tensile strength is generally low (<480MPa). Among them, the technology announced by patent CN10358997A can prepare alloys with a tensile strength exceeding 480MPa, but the preparation process is cumbersome and the organizational stability is difficult to guarantee. In addition, patents CN105506521 and CN108149172B announce the processing method of brass texture or fine-grained cubic texture fatigue-resistant aluminum alloy sheet, and the ultimate tensile strength of the alloy is only about 400MPa. In the currently published patent technologies, the alloy yield ratio is relatively high. When the alloy yield ratio increases, the plastic damage space is reduced, which means that dislocation slip will be inhibited, increasing the stress concentration in the fatigue process, further increasing the tendency of microcrack formation, thereby reducing fatigue damage resistance. It can be seen that it is very difficult to make the alloy take into account both high strength and toughness and excellent fatigue damage resistance. The existing published patent technologies are difficult to solve the problem of Al-Cu-Mg-based alloys having excellent fatigue damage resistance when the ultimate tensile strength exceeds 500MPa.

[0004] In summary, the 2000 series aviation aluminum alloys disclosed in the prior art are very difficult to achieve high strength and toughness, excellent fatigue damage resistance and low production cost. With the development of the aviation industry, the alloy performance indicators in the currently published patented technologies are difficult to meet the design requirements of future aviation aircraft. Summary of the invention

[0005] In order to solve the problem that it is difficult for aluminum alloys to have both high strength and toughness and excellent fatigue damage resistance, the present invention provides a high-strength aluminum alloy with excellent fatigue damage resistance. The high-strength aluminum alloy has both high strength and toughness and excellent fatigue damage resistance, has a reasonable material composition design, a simple preparation process flow, and low industrial production cost.

[0006] The present invention also provides a method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance.

[0007] The present invention is achieved through the following technical solutions:

[0008] The present invention provides a high-strength aluminum alloy with excellent fatigue damage resistance. The chemical composition of the high-strength aluminum alloy is as follows, measured by 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 comprises Goss+Cube or Goss+P grain bands, the major diameter 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.

[0011] Preferably, 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%;

[0012] 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 / cyc le, instability ΔK can reach 39MPa*m 1 / 2 above.

[0013] Based on the same inventive concept, the present invention provides a method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance, the preparation method comprising:

[0014] An ingot is obtained by a smelting and casting process;

[0015] The ingot is subjected to a primary homogenization treatment, and then the temperature is raised to perform a secondary homogenization treatment to obtain a homogenized ingot;

[0016] The homogenized ingot is subjected to large deformation hot rolling to obtain a hot-rolled plate;

[0017] Alternatively, the homogenized ingot is subjected to large deformation hot rolling, and then cold rolled or directly cold rolled after annealing to obtain a cold rolled sheet;

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

[0019] 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;

[0020] Wherein, the chemical composition of the ingot is as follows, measured by 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 is Al, 3≤Cu / Mg≤4.

[0022] Preferably, the ingot is subjected to a primary homogenization treatment, and then the temperature is raised to perform a secondary homogenization treatment to obtain a homogenized ingot, which specifically includes:

[0023] The ingot is subjected to a primary homogenization treatment at 465-490° C. for 24-48 hours;

[0024] After the primary homogenization treatment is completed, the temperature is raised to 490°C to 505°C for secondary homogenization treatment. The homogenization time is 12 to 48 hours to obtain a homogenized ingot.

[0025] Preferably, the step of subjecting the homogenized ingot to large deformation hot rolling to obtain a hot-rolled plate specifically comprises:

[0026] The homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm.

[0027] Preferably, the homogenized ingot is subjected to large deformation hot rolling, and then cold rolled or directly cold rolled after annealing to obtain a cold rolled sheet, specifically comprising:

[0028] The homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm, and then the hot-rolled plate is directly cold-rolled with a cold rolling deformation of ≤50%, to obtain a cold-rolled plate;

[0029] Alternatively, the homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm, the hot-rolled plate is annealed at a temperature of 260-330° C. for 2-4 h, and the annealed plate is subsequently cold rolled with a cold rolling deformation of ≤50% to obtain a cold-rolled plate.

[0030] Preferably, the step of isothermally annealing the hot rolled sheet or the cold rolled sheet to obtain the annealed aluminum alloy sheet specifically comprises:

[0031] Isothermally annealing the hot-rolled sheet or the cold-rolled sheet to obtain an annealed aluminum alloy sheet;

[0032] Among them, when the large deformation hot rolling temperature is lower than 450°C, the isothermal annealing temperature is controlled at 320°C ~ 370°C, and the annealing time is 2 to 5 hours. When the large deformation hot rolling temperature is higher than 450°C, the isothermal annealing temperature is controlled at 340°C ~ 420°C, and the annealing time is 3 to 6 hours.

[0033] Preferably, 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, specifically comprising:

[0034] The annealed aluminum alloy plate is subjected to high temperature short-time solution treatment, the solution treatment temperature is 475-505° C., the time is 3-30 minutes, and then water quenching is performed with water as the medium, the quenching transfer time is ≤15 seconds, and then aging treatment is performed to obtain a high-strength aluminum alloy;

[0035] Wherein, when the aging treatment adopts the natural aging system, the natural aging time is ≥240h, and when the artificial aging system is adopted, the artificial aging temperature is 140°C to 180°C, and the aging time is 1 to 5h.

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

[0037] One or more technical solutions in the embodiments of the present invention have at least the following technical effects or advantages:

[0038] 1. The present invention discloses a high-strength aluminum alloy with excellent fatigue damage resistance. The high-strength aluminum alloy solves the problem that the existing Al-Cu-Mg-based alloys have both high strength and toughness and excellent fatigue damage resistance, and caters to the concept of low-cost preparation. The aged high-strength aluminum alloy prepared by the present invention has a tensile strength of up to 500MPa and an elongation after fracture of 18-30%. Its strength and toughness are very excellent. The fatigue crack growth rate of the aged alloy plate in the TL direction is: ΔK=20MPa*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 / cyc le, instability ΔK can reach 39MPa*m 1 / 2 The high-strength aluminum alloy material of the present invention has a reasonable composition design and a simple preparation process, which meets the design requirements of contemporary and future aviation vehicles and is suitable for large-scale industrial production and application.

[0039] 2. The present invention discloses a method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance. The present invention controls the texture and microstructure of the aged Al-Cu-Mg based alloy by deformation and microalloying, so that the high-strength aluminum alloy can obtain an ultimate tensile strength of more than 500MPa while also obtaining very excellent fatigue damage resistance. The high-strength aluminum alloy prepared by this method can maintain high strength and toughness and excellent fatigue damage resistance in the aged state even after different deformations and heat treatments. The present invention solves the problem that the current Al-Cu-Mg based alloy is difficult to have both high strength and toughness, excellent fatigue damage resistance and low-cost preparation, and meets the design indicators of current and future aviation vehicles. BRIEF DESCRIPTION OF THE DRAWINGS

[0040] In order to more clearly illustrate the technical solutions in the embodiments of the present invention, the following briefly introduces the drawings required for use in the description of the embodiments. Obviously, the drawings described below are some embodiments of the present invention. For ordinary technicians in this field, other drawings can be obtained based on these drawings without creative work.

[0041] Figure 1 It is the orientation distribution function ODF diagram of the hot-rolled alloy plate of Example 1;

[0042] Figure 2This is the orientation distribution function ODF diagram of the naturally aged plate in Example 1;

[0043] Figure 3 is the fatigue crack growth rate curve of the alloy in Example 1;

[0044] Figure 4 This is the natural aging TEM structure of the alloy in Example 1;

[0045] Figure 5 It is the orientation distribution function ODF diagram of the hot-rolled alloy plate of Example 2;

[0046] Figure 6 This is the orientation distribution function ODF diagram of the natural aging state plate in Example 2;

[0047] Figure 7 is the fatigue crack growth rate curve of the alloy in Example 2;

[0048] Figure 8 This is the natural aging TEM structure of the alloy in Example 2;

[0049] Fig. 9 is the fatigue crack growth rate curve of the alloy in Example 3;

[0050] Fig.10 The ODF diagram of the orientation distribution function of the plate in the direct cold-rolled state after hot-rolling of the alloy in Example 4;

[0051] Fig.11 This is the orientation distribution function ODF diagram of the naturally aged plate in Example 4;

[0052] Fig.12 The fatigue crack growth rate curve and tensile and fatigue performance indicators of the alloy of Example 4;

[0053] Fig.13 TEM structure of the alloy in the natural aging state of Example 4;

[0054] Fig.14 This is a schematic diagram of the grain structure distribution of the naturally aged plate of Example 4;

[0055] Fig.15 This is the orientation distribution function ODF diagram of the naturally aged plate in Example 5;

[0056] Fig.16 is the fatigue crack growth rate curve of the alloy in Example 5;

[0057] Fig.17 TEM structure of the alloy in the natural aging state of Example 6;

[0058] Fig.18 is the fatigue crack growth rate curve of the alloy in Example 6;

[0059] Fig.19 This is the artificial aging TEM structure of the alloy in Example 6;

[0060] Fig. 20 is the fatigue crack growth rate curve of the alloy of Example 7;

[0061] Fig.21 This is the artificial aging TEM structure of the alloy in Example 7;

[0062] Fig. 22 It is the orientation distribution function ODF diagram of the hot-rolled alloy plate of comparative example 1;

[0063] Fig.23 It is the orientation distribution function ODF diagram of the natural aging state plate of comparative example 1;

[0064] Fig.24 is the fatigue crack growth rate curve of the alloy of comparative example 1;

[0065] Fig.25 TEM structure of the alloy in the natural aging state of Comparative Example 1;

[0066] Fig.26 It is the orientation distribution function ODF diagram of the hot-rolled sheet of alloy in comparative example 2;

[0067] Fig. 27 It is the orientation distribution function ODF diagram of the natural aging state plate of comparative example 2;

[0068] Fig.28 is the fatigue crack growth rate curve of the alloy of comparative example 2;

[0069] Fig.29 TEM structure of the alloy in the natural aging state of Comparative Example 2;

[0070] Fig.30 is the fatigue crack growth rate curve of the alloy of comparative example 3;

[0071] Fig.31 This is the artificial aging TEM structure of the alloy in comparative example 3;

[0072] Fig.32 It is the orientation distribution function ODF diagram of the hot-rolled plate of the alloy in comparative example 4;

[0073] Fig.33 It is the orientation distribution function ODF diagram of the natural aging state plate of comparative example 4;

[0074] Fig.34 is the fatigue crack growth rate curve of the alloy of comparative example 4;

[0075] Fig.35This is the natural aging TEM structure of the alloy in comparative example 4. DETAILED DESCRIPTION

[0076] The present invention will be described in detail below in conjunction with specific implementations and examples, and the advantages and various effects of the present invention will be more clearly presented. It should be understood by those skilled in the art that these specific implementations and examples are used to illustrate the present invention, rather than to limit the present invention.

[0077] Throughout the specification, unless otherwise specifically stated, the terms used herein should be understood as meanings commonly used in the art. Therefore, unless otherwise defined, all technical and scientific terms used herein have the same meanings as those generally understood by those skilled in the art to which the present invention belongs. In the event of a conflict, the present specification takes precedence.

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

[0079] The technical principles of the present invention are as follows:

[0080] The present invention provides a method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance, the preparation method comprising:

[0081] S1. Obtaining an ingot by a smelting and casting process;

[0082] S2. The ingot is subjected to a primary homogenization treatment, and then the temperature is raised for a secondary homogenization treatment to obtain a homogenized ingot;

[0083] S3. The homogenized ingot is subjected to large deformation hot rolling to obtain a hot-rolled plate;

[0084] Alternatively, the homogenized ingot is subjected to large deformation hot rolling, and then cold rolled or directly cold rolled after annealing to obtain a cold rolled sheet;

[0085] S4. Isothermally annealing the hot-rolled sheet or the cold-rolled sheet to obtain an annealed aluminum alloy sheet;

[0086] S5. 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;

[0087] Wherein, the chemical composition of the ingot is as follows, measured by mass fraction:

[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 invention, the beneficial effects of using the above chemical elements and proportions in the ingot are as follows:

[0090] Cu: 4.5-5.5%. Within this range, the solid solution strengthening provided by the Cu element and the cluster strengthening provided by the combination of Cu and Mg elements to form Cu-Mg atomic clusters can be maximized. Below this mass fraction range, the solid solution strengthening provided by the Cu element is significantly weakened. Above this mass fraction range, the Cu element will form too much coarse Al2Cu phase with Al. Since the solid solubility limit of Cu in Al is reached, it is difficult to completely dissolve into the matrix after homogenization treatment or subsequent solid solution treatment. The coarse Al2Cu phase will not only weaken the strength, but also significantly hinder the reciprocating sliding of dislocations during fatigue, greatly increasing the fatigue crack growth rate.

[0091] Mg: 1.0-1.4%. Within this range, the solid solution strengthening provided by the Mg element and the cluster strengthening effect provided by the combination of Mg and Cu elements to form Cu-Mg atomic clusters in the aging state can be maximized. The design of the Mg element matches the Cu element. When the mass fraction of the Cu element is determined, if the Mg element content is lower than this range, the solid solution strengthening and cluster strengthening will be weakened. If it is higher than this range, the Al2CuMg phase content formed by Mg, Cu and Al in the cast state is high, which greatly increases the difficulty of homogenization and the risk of overburning.

[0092] Ag: 0.2-1%. Ag addition can provide certain solid solution strengthening and contribute to precipitation strengthening. Within this mass fraction range, the Ag element can significantly reduce the dislocation content formed by the collapse of supersaturated vacancies in the natural aging state, which is beneficial to the reciprocating slip of dislocations during fatigue and increases the resistance to fatigue crack growth. 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 moderate-sized Cu-Mg atomic clusters not only provide precipitation strengthening, but also are not easy to cause cyclic softening effects, ensuring high strength while also having excellent fatigue properties. Below this mass fraction range, the ability to suppress dislocation formation is poor in the natural aging state, and the artificial aging state cannot effectively promote the formation of clusters. Above this mass fraction range, the material preparation cost will increase significantly.

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

[0094] Ti: 0.06-0.10%. The addition of Ti is mainly for solid solution strengthening, and secondly, it will produce a certain grain refinement effect. Within this range, a better strengthening effect can be guaranteed, and the content of coarse phase formed by Ti, Al and Si can be reduced, as well as the material preparation cost can be reduced.

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

[0096] Zn: <0.1%, Zn is an impurity element introduced during the smelting process. Controlling Zn <0.1% can effectively reduce the consumption of Mg by Zn. If the Zn content is too high, Zn atoms will migrate to the grain boundaries during the recrystallization process, inhibiting the formation of a strong Goss texture.

[0097] Among them, the advantage of 3≤Cu / Mg≤4 is that within this Cu / Mg ratio range, the number density, size and distribution of coarse Al2CuMg and Al2Cu second phases that are supersaturated and precipitated in the cast alloy can be effectively controlled, and these phases are easily dissolved into the matrix after homogenization treatment. The reduction in the content of these coarse second phases, on the one hand, means that solid solution strengthening and precipitation strengthening can be maximized, and fatigue crack sources are also reduced. On the other hand, the reduction in the content of these coarse phases is conducive 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 properties of the alloy.

[0098] Preferably, step S2 specifically includes:

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

[0100] After the primary homogenization treatment is completed, the temperature is raised to 490°C to 505°C for secondary homogenization treatment. The homogenization time is 12 to 48 hours to obtain a homogenized ingot.

[0101] In the present invention, the advantage of the primary homogenization treatment temperature being 465-490°C is that within this temperature range, the Al2CuMg phase in the cast alloy can be dissolved into the matrix as much as possible while ensuring that these phases are not overburned. Below this temperature range, the Al2CuMg phase is difficult to be completely eliminated, and above this temperature range, the alloy is overburned.

[0102] In the present invention, the advantage of the secondary homogenization treatment temperature being 490°C to 505°C is that it can ensure that the Al2Cu phase in the 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 invention, the purpose of the isothermal (i.e., there is no variable temperature heating during the homogenization process, but heat treatment at the same temperature, and the variable is time) two-stage homogenization treatment is to dissolve the coarse Al2CuMg and Al2Cu phases supersaturated and precipitated in the alloy into the matrix to the maximum extent.

[0104] Preferably, step S3 specifically includes:

[0105] The homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm.

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

[0107] In the present invention, the large deformation hot rolling temperature is 390-490°C. Within this temperature range, the sheet forming effect is optimal and high-strength copper and brass deformation textures can be obtained. Too low a hot rolling temperature will cause serious edge cracks in the sheet, and too high a hot rolling temperature will increase the risk of overburning and cracking of the sheet. In addition, the expected texture components cannot be obtained. The advantage of a total rolling deformation greater than 85% is that high-strength copper and brass textures can be obtained.

[0108] Furthermore, before the hot-rolled plate is cold-rolled, the hot-rolled plate is annealed at a temperature of 260° C. to 330° C. and for a time of 2 to 4 hours.

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

[0110] Preferably, step S4 specifically includes:

[0111] Isothermally annealing the hot-rolled sheet or the cold-rolled sheet to obtain an annealed aluminum alloy sheet;

[0112] Among them, when the large deformation hot rolling temperature is lower than 450°C, the isothermal annealing temperature is controlled at 320°C ~ 370°C, and the annealing time is 2 to 5 hours. When the large deformation hot rolling temperature is higher than 450°C, the isothermal annealing temperature is controlled at 340°C ~ 420°C, and the annealing time is 3 to 6 hours.

[0113] In the present invention, when the large deformation hot rolling temperature is lower than 450°C, the isothermal annealing temperature is controlled at 320°C to 370°C, and the annealing time is 2 to 5 hours. The advantage is that high-strength Gsoss and P textures can be obtained. When the large deformation hot rolling temperature is higher than 450°C, the isothermal annealing temperature is controlled at 340°C to 420°C, and the annealing time is 3 to 6 hours. The advantage is that high-strength Goss and Cube textures can be obtained at this temperature and time.

[0114] Preferably, step S5 specifically includes:

[0115] The annealed aluminum alloy plate is subjected to high temperature short-time solution treatment, the solution treatment temperature is 475-505° C., the time is 3-30 minutes, and then water quenching is performed with water as the medium, the quenching transfer time is ≤15 seconds, and then aging treatment is performed to obtain a high-strength aluminum alloy;

[0116] Wherein, when the aging treatment adopts the natural aging system, the natural aging time is ≥240h, and when the artificial aging system is adopted, the artificial aging temperature is 140°C to 180°C, and the aging time is 1 to 5h.

[0117] In the present invention, the solution treatment temperature is 475-505° C. and the time is 3-30 min. The advantage is that within the temperature and time range, the strength of the Goss, P, and Cube textures formed by isothermal annealing will not be reduced, and the fine dispersed S phase particles in the matrix can be fully dissolved into the matrix, thereby ensuring the solid solution and precipitation strengthening effects. The water quenching is carried out with water as the medium, and the quenching transfer time is ≤15 s. The advantage is that the alloy is in a high solid solubility, which is beneficial to promote the precipitation of Cu-Mg atomic clusters during the aging process.

[0118] In the present invention, when the natural aging system is adopted, the advantage of the natural aging time being ≥240h is that the solid solution atoms can be fully precipitated to produce high-density Cu-Mg atomic clusters, and the alloy is in a relatively stable state. When the artificial aging system is adopted, the artificial aging temperature is 140°C to 180°C, and the aging time is 1 to 5h. The advantage is that Cu-Mg atomic clusters with high number density and moderate size can be formed within this temperature and time range.

[0119] The yield strength of the aged aluminum alloy sheet prepared by the above technical route is 300-350MPa, the tensile strength is 495-520MPa, and the elongation after fracture is 18-30%. The fatigue damage index of the aged alloy in the TL direction is: ΔK = 20MPa*m1 / 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 / cyc le, instability ΔK can reach 39MPa*m 1 / 2 The aged alloy is controlled by a strong Goss+Cube or Goss+P texture, the combined strength of the two textures is >10, the volume fraction is >20%, and the strength of the Goss and P texture components in the combination is >7. The matrix contains a certain number of Goss+Cube or Goss+P grain bands, the long diameter of the grains inside the grain bands is parallel to the TD direction, the grain boundaries are perpendicular to the RD direction, and these grain bands are surrounded by equiaxed grains.

[0120] The inventors found that by regulating the stacking fault energy of the Al-Cu-Mg alloy by the Cu / Mg ratio and regulating the subsequent annealing process, a combination of Goss+P textures can be formed. When these two texture components have a large torsion or deflection angle with the surrounding grains, the cracks usually deflect significantly, and the {111} planes of these two components are close to the direction of the maximum critical shear stress, which alleviates the stress concentration during crack propagation. Therefore, the crack propagation resistance is increased, thereby improving fatigue performance. However, although CN103045976A, CN103526140A, CN10358997A, CN108103373B and CN108504915B published by the inventors in the early stage have excellent fatigue performance, it is difficult to achieve a material ultimate tensile strength of more than 480MPa. Generally, high 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. Although high-density precipitated phases and fine grains are beneficial to improving strength, these phases will hinder the reciprocating motion of dislocations during fatigue, thereby reducing fatigue performance. Furthermore, the fatigue crack growth rate detected in the above-mentioned patent is in the LT direction, and the long diameter direction of Goss and other grains is along the RD, which also means that the grain boundaries are mostly parallel to the RD, that is, the crack growth direction. Usually the grain boundary is a weak area, and the grain boundary bonding force is weak, which means that the crack growth 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 (>495MPa) and have a lower fatigue crack growth rate in the TL direction.

[0121] The present invention controls the texture and microstructure of the aged Al-Cu-Mg-Ag alloy by both deformation and microalloying, so that the alloy of the invention can obtain an ultimate tensile strength of more than 500MPa and also obtain very excellent fatigue damage resistance. The principle of the present invention is to control the texture, grain structure, coarse second phase particles and reduce dislocation density of the aged alloy by controlling the deformation temperature and microalloying. The specific principle is that under large deformation, deformation at different temperatures can make the alloy form deformed grains with a large aspect ratio extending along the RD direction. Under the influence of the deformation temperature, the content of Brass and Copper textures in the deformed matrix can be regulated. In the subsequent isothermal recrystallization annealing process, since the recrystallized grain growth originates from the deformed grains, and the grain boundaries act as barriers to limit the expansion of the recrystallized grains in the TD direction, 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, it eventually leads to the formation of Goss or Goss+P grain bands with the long diameter direction of the grains along the TD direction. A nearly equiaxed crystal structure is formed around it. Therefore, Goss or Goss+P grain bands are formed in the aged alloy. The grain growth direction inside 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-state propagation stage, it is difficult for the crack to pass through these grains, but it propagates around the grain boundaries of these grains, which greatly increases the crack propagation range. And to a certain extent, the formation of surface slip bands is inhibited, the crack initiation caused by the immersion of slip bands 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, the number density and distribution of coarse Al2CuMg and Al2Cu second phases in the alloy are controlled by deformation process and microalloying. Reducing the number density of these two coarse phases promotes the formation of fatigue striation bands in the steady-state extension zone, and dislocation reciprocating slip becomes easier, thereby inhibiting the generation of microcracks between striation bands. Adding Ag element on the basis of controlling the Cu / Mg ratio, due to the capture of vacancies by Ag, ultimately reduces the content of ordered dislocations in the aged alloy, further reducing the dislocation slip resistance and accelerated expansion of stress concentration during fatigue crack propagation. In addition, at a high Cu / Mg, the Cu atoms in the excess phase after homogenization are dissolved into the matrix to the maximum extent, eliminating dendrite segregation, and producing a solid solution strengthening effect in the subsequent solid solution treatment. The reduction in the content of coarse Al2CuMg and Al2Cu phases inhibits the cracking tendency during plastic deformation to a certain extent, produces continuous hardening ability, and thus contributes to higher tensile strength. Furthermore, the grain boundaries in the Goss or Goss+P grain bands are perpendicular to the direction of tensile stress, which further increases the deformation resistance. The improvement of plastic damage space, that is, the yield strength ratio, lies in the secondary recrystallization and the formation of atomic clusters suppressed by short-term solid solution, thereby obtaining a low yield strength ratio.In addition, by controlling the Zn content, the consumption of Zn on the Mg element is reduced, and Zn atoms migrate along the grain boundaries during the recrystallization process, inhibiting the formation of a strong Goss texture. With the support of the above principles and matching the technology disclosed in the present invention, the Al-Cu-Mg-Ag alloy can obtain an ultimate tensile strength of more than 500MPa, while the elongation after fracture is guaranteed to be between 18% and 30%, and has excellent fatigue damage resistance.

[0122] The following will describe in detail a high-strength aluminum alloy with excellent fatigue damage resistance and a preparation method thereof in combination with embodiments and experimental data.

[0123] Example 1

[0124] The chemical composition of the high-strength aluminum alloy in this embodiment 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 is Al. First, the alloy composition is divided and smelted into ingots (after smelting, the alloy elements will fluctuate, Cu and Mg elements fluctuate within ±0.2%, Ag and Mn elements fluctuate within ±0.1%, and Ti element fluctuates within ±0.02%, and Fe, Si and Zn elements are controlled within the designed composition range). Subsequently, the ingot is homogenized at 470℃ / 48h+505℃ / 12h, and then hot rolled at 405℃ with a deformation of 95.6%, and the plate thickness is 2mm. This was followed by isothermal annealing at 340 °C for 5 h, followed by solution treatment at 485 °C for 30 min, followed by immediate water quenching with a quenching transfer time of 5 s and natural aging for 240 h.

[0125] The hot rolled alloy is controlled by Copper texture, with Brass texture strength of 6.82, Copper texture strength of 8.20, and S texture strength of 6.82. The naturally aged alloy is controlled by Goss+P texture, with Goss texture strength of 8.01 and P texture strength of 8.01. Figure 1 and 2 Ph i 1, PHI, Constant Ph i2 represent the φ1, Φ, φ2 angles in the Euler space. The dislocation content in the naturally aged alloy is low. The yield strength of the naturally aged alloy is 299.9MPa, the tensile strength is 496.3MPa, the elongation after fracture is 22.5%, and the yield strength ratio is 0.604. The fatigue crack growth rate in the TL direction is shown in Table 1, and the instability ΔK = 30.2MPa*m 1 / 2 .

[0126] Attached Figure 3The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, mm / cyc le represents the distance of crack growth per cycle, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor. The room temperature tensile test was carried out at a rate of 2mm / min on an MTS Landmark universal testing machine according to GB228-87 "Metal Tensile Test Methods". The yield strength was obtained by clamping an extensometer on the sample, the tensile strength was directly taken from the parameter displayed by the equipment, and the elongation after fracture was measured with a vernier caliper. The fatigue crack growth rate was carried out on an MTS 810 fatigue testing machine according to GB / T63982000. Before the test, the sample was polished with 180#, 600#, and 1000# grit sandpaper until there were no defects on the sample surface. The polishing process ensured the uniformity of the sample thickness. The test samples were all standard compact tensile (C(T)) specimens with a thickness of 2±0.2mm, and the samples were taken along the TL direction of the rolled plate. The test methods / testing standards / testing equipment for yield strength, tensile strength, elongation after fracture and fatigue crack growth rate in the TL direction in the subsequent examples and comparative examples are consistent with those in Example 1.

[0127] Table 1 Fatigue crack growth rate (×10 -4 mm / cycle)

[0128]

[0129] Example 2

[0130] The chemical composition of the high-strength aluminum alloy in this embodiment 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 composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuations of Ti elements are within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Then the ingot is homogenized at 475℃ / 36h+500℃ / 24h, and then hot rolled at 490℃ with a deformation of 90%, and the plate thickness is 3mm. Then isothermal annealing is performed at 390℃ for 4h, and then solution treatment is performed at 495℃ for 15min, followed by immediate water quenching, quenching transfer time of 3s, and natural aging for 260h.

[0131] The hot-rolled alloy is controlled by Brass texture, with a Brass texture strength of 8.68, a Copper texture strength of 4.32, and an S texture strength of 4.32. The naturally aged alloy is controlled by Goss+Cube texture, with a Goss texture strength of 9.28 and a Cube texture strength of 9.28. Figure 5 and 6 Ph i 1, PHI, Constant Ph i2 represent the φ1, φ, φ2 angles in the Euler space. There are locally ordered coil dislocations in the naturally aged alloy with low density. The yield strength of the naturally aged alloy is 311.6MPa, the tensile strength is 507.4MPa, the elongation after fracture is 23%, and the yield strength ratio is 0.61. The fatigue crack growth rate in the TL direction is shown in Table 2, and the instability ΔK = 30.4MPa*m 1 / 2 Attached Figure 7 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0132] Table 2 Fatigue crack growth rate of Example 2 alloy at different ΔK (×10 -4 mm / cycle)

[0133]

[0134] Example 3

[0135] The chemical composition of the high-strength aluminum alloy in this embodiment 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 composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuation of Ti element is within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Subsequently, the ingot is homogenized at 465℃ / 36h+495℃ / 48h, and then hot rolled at 435℃ with a deformation of 93%, and the thickness of the hot rolled plate is 2mm. This was followed by isothermal annealing at 360 °C for 5 h, followed by solution treatment at 503 °C for 3 min, followed by immediate water quenching with a quenching transfer time of 5 s and natural aging for 280 h.

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

[0137] Table 3 Fatigue crack growth rate of Example 3 alloy at different ΔK (×10-4 mm / cycle)

[0138]

[0139] Example 4

[0140] The chemical composition of the high-strength aluminum alloy in this embodiment is Cu 5.05%, Mg 1.3%, Ag 0.65%, Mn 0.82%, Ti0.06%, Fe 0.05%, Si 0.04%, Zn 0.07%, and the balance is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuation of Ti element is within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Then the ingot is homogenized at 480℃ / 36h+495℃ / 24h, and then hot rolled at 455℃ with a deformation of 85%, and then directly cold rolled at 45%, and the thickness of the cold rolled plate is 1.8mm. This was followed by isothermal annealing at 365 °C for 6 h, followed by solution treatment at 500 °C for 10 min, followed by immediate water quenching with a quenching transfer time of 7 s and natural aging for 280 h.

[0141] The cold-rolled alloy is controlled by Brass texture, with a Brass texture strength of 9.64, a Copper texture strength of 3.18, and an S texture strength of 3.18. The naturally aged alloy is controlled by Goss+Cube texture, with a Goss texture strength of 7.1 and a Cube texture strength of 4.1. Fig.10 and 11 Ph i 1, PHI, Constant Ph i2 represent the φ1, Φ, φ2 angles in the Euler space. Only a small amount of ordered dislocation loops and coil dislocations exist in the naturally aged alloy. The microstructure contains Goss+Cube grain bands. The long diameter 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 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 growth rate in the TL direction is shown in Table 4, and the instability ΔK=37.7MPa*m 1 / 2 Attached Fig.12 The vertical axis da / dN (mm / cycle) represents the crack growth rate, and the horizontal axis DeltaK (MPa*m1 / 2) represents the stress intensity factor.

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

[0143]

[0144] Example 5

[0145] The chemical composition of the high-strength aluminum alloy in this embodiment 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%, and the balance is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuation of Ti element is within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Then the ingot is homogenized at 485℃ / 24h+495℃ / 36h, and then hot rolled at 465℃ with a deformation of 91%. After the hot-rolled plate is annealed at 300℃ / 3h, it is then cold rolled by 30%, and the thickness of the cold-rolled plate is 2mm. This was followed by isothermal annealing at 370 °C for 4 h, followed by solution treatment at 505 °C for 6 min, followed by immediate water quenching, a quenching transfer time of 10 s, and natural aging for 265 h.

[0146] The naturally aged alloy is controlled by Goss+Cube texture, with Goss texture strength of 10.7 and Cube texture strength of 8.2. Fig.15 Ph i 1, PHI, Constant Ph i2 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 coil dislocations in the naturally aged alloy. The yield strength of the naturally aged 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 growth rate in the TL direction is shown in Table 5. Instability ΔK = 39.4MPa*m 1 / 2 Attached Fig.16 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate Del taK (MPa*m1 / 2) represents the stress intensity factor.

[0147] Table 5 Fatigue crack growth rate of Example 5 alloy at different ΔK (×10 -4 mm / cycle)

[0148]

[0149] Example 6

[0150] The chemical composition of the high-strength aluminum alloy in this embodiment 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 is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuation of Ti element is within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Subsequently, the ingot is homogenized at 490℃ / 36h+490℃ / 36h, and then hot rolled at 480℃ with a deformation of 91.5%, and the plate thickness is 2.5mm. Then it was isothermal annealed at 415℃ for 3h, then solution treated at 495℃ for 8min, and then immediately water quenched, with a quenching transfer time of 11s, and an artificial aging temperature of 155℃ for 3h.

[0151] Dislocations in the artificially aged alloy almost disappear, and low-density atomic clusters exist. The yield strength of the artificially aged 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 growth rate in the TL direction is shown in Table 6, and the instability ΔK = 35.1MPa*m 1 / 2 Attached Fig.18 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0152] Table 6 Fatigue crack growth rate of Example 6 alloy at different ΔK (×10 -4 mm / cycle)

[0153]

[0154] Example 7

[0155] The chemical composition of the high-strength aluminum alloy in this embodiment is Cu 4.78%, Mg 1.4%, Ag 0.6%, Mn 0.77%, Ti0.083%, Fe 0.032%, Si 0.034%, Zn 0.043%, and the balance is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuation of Ti element is within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Subsequently, the ingot is homogenized at 485℃ / 24h+490℃ / 48h, and then hot rolled at 470℃ with a deformation of 87%, and the plate thickness is 3.2mm. Then it was isothermal annealed at 400℃ for 4h, then solution treated at 490℃ for 20min, and then immediately water quenched, with a quenching transfer time of 6s, an artificial aging temperature of 173℃, and an aging time of 1h.

[0156] Dislocations in the artificially aged alloy almost disappeared, and low-density atomic clusters existed. Compared with Example 5, the size of the atomic clusters increased slightly. The yield strength of the artificially aged alloy was 341.4 MPa, the tensile strength was 518.6 MPa, the elongation after fracture was 23.4%, and the yield strength ratio was 0.658. The fatigue crack growth rate in the TL direction is shown in Table 7, and the instability ΔK = 31.2 MPa*m 1 / 2 Attached Fig. 20 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

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

[0158]

[0159] Comparative Example 1

[0160] Comparative Example 1 The chemical composition is Cu 4.1%, Mg 1.7%, Mn 0.71%, Ti 0.07%, Fe 0.045%, Si 0.032%, Zn 0.06%, and the balance is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, the fluctuations of Mn elements are within ±0.1%, and the fluctuations of Ti elements are within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Then the ingot is homogenized at 475℃ / 36h+500℃ / 24h, and then hot rolled at 490℃ with a deformation of 90%, and the plate thickness is 3mm. Then isothermal annealing is performed at 390℃ for 4h, and then solution treatment is performed at 495℃ for 15min, followed by immediate water quenching, quenching transfer time of 3s, and natural aging for 260h.

[0161] The hot-rolled plate 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 the Goss+Cube texture, with a Goss texture strength of 7.87 and a Cube texture strength of 7.87. Fig. 22 and 23 Ph i 1, PHI, Constant Ph i2 represent the φ1, Φ, φ2 angles in the Euler space. The naturally aged alloy is composed of highly ordered dislocation rings with a high dislocation density. The yield strength of the naturally aged alloy is 305.7MPa, the tensile strength is 477.5MPa, the elongation after fracture is 23.6%, and the yield strength ratio is 0.64. It can be seen that although the fatigue performance of the naturally aged alloy in Comparative Example 1 is good, the ultimate tensile strength is far from 500MPa. The fatigue crack growth rate in the TL direction is shown in Table 8, and the instability ΔK=33.5MPa*m 1 / 2 Attached Fig.24 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0162] Table 8 Fatigue crack growth rate of alloy of comparative example 1 at different ΔK (×10 -4 mm / cycle)

[0163]

[0164]

[0165] Comparative Example 2

[0166] The chemical composition of the alloy of comparative example 2 is Cu 4.83%, Mg 1.25%, Mn 0.6%, Ti 0.08%, Fe0.041%, Si 0.036%, Zn 0.08%, and the balance is Al. The alloy has a high Cu / Mg ratio, but no Ag element is added. First, the alloy is divided according to the above alloy composition and then smelted into an ingot. First, the alloy is divided according to the above alloy composition and then smelted into an ingot. After smelting, the alloy elements will fluctuate, and the fluctuations of Cu and Mg elements are within ±0.2%, the fluctuations of Mn elements are within ±0.1%, and the fluctuations of Ti elements are within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Subsequently, the ingot is homogenized at 475℃ / 34h+500℃ / 20h, and then hot rolled at 486℃ with a deformation of 91%, and the plate thickness is 2.8mm. This was followed by isothermal annealing at 385 °C for 4 h, followed by solution treatment at 495 °C for 15 min, followed by immediate water quenching with a quenching transfer time of 4 s and natural aging for 260 h.

[0167] The hot-rolled plate 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 the Goss+Cube texture, with a Goss texture strength of 9.96 and a Cube texture strength of 9.96. Fig.26 and 27 Ph i 1, PHI, Constant Ph i2 represent the φ1, Φ, φ2 angles in the Euler space. The naturally aged alloy is composed of highly ordered dislocation rings with a high dislocation density. The yield strength of the naturally aged alloy is 311.9MPa, the tensile strength is 494.8MPa, the elongation after fracture is 21.9%, and the yield strength ratio is 0.63. It can be seen that the fatigue performance of the naturally aged alloy in Comparative Example 2 is poor, especially under large ΔK, the crack growth rate is significantly accelerated, and the tensile strength has not yet exceeded 500MPa. ΔK = 30MPa*m 1 / 2 , the crack growth rate is as high as 3.5×10 -3 mm / cycle. The fatigue crack growth rate in the TL direction is shown in Table 9. The instability ΔK = 30.1 MPa*m 1 / 2 Attached Fig.28 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0168] Table 9 Fatigue crack growth rate of alloy of comparative example 2 at different ΔK (×10 -4 mm / cycle)

[0169]

[0170] Comparative Example 3

[0171] The chemical composition of the high-strength aluminum alloy in 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 is Al. First, the above alloy components are mixed and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, and the fluctuations of Mn elements are within ±0.1%.

[0172] The Ti element fluctuates within ±0.02%, and the Fe, Si and Zn elements are controlled within the designed composition range. The ingot is then homogenized at 490℃ / 35h+490℃ / 36h, and then hot rolled at 480℃ with a deformation of 91%, and the plate thickness is 2.4mm. It is then isothermally annealed at 410℃ for 3h, then solution treated at 495℃ for 7min, followed by immediate water quenching, quenching transfer time of 10s, artificial aging temperature of 159℃, and aging time of 3h.

[0173] The artificially aged alloy is also composed of highly ordered dislocation rings with a high dislocation density. The yield strength of the artificially aged 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 alloy in Example 3 drops sharply, and the ultimate tensile strength is far from 500MPa. Although the crack growth rate is slow under low ΔK, ΔK = 26.6MPa*m 1 / 2 The fatigue crack growth rate in the TL direction is shown in Table 10. Fig.30 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0174] Table 10 Fatigue crack growth rate of alloy of comparative example 3 at different ΔK (×10 -4 mm / cycle)

[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%, Ti0.08%, Fe 0.045%, Si 0.045%, Zn 0.09%, and the balance is Al. First, the alloy composition is divided and smelted into ingots. After smelting, the alloy elements will fluctuate. The fluctuations of Cu and Mg elements are within ±0.2%, Ag and Mn elements are within ±0.1%, and the fluctuations of Ti elements are within ±0.02%. Fe, Si and Zn elements are controlled within the designed composition range. Then the ingot is homogenized at 490℃ / 36h+495℃ / 24h, and then hot rolled at 480℃ with a deformation of 90%, and the plate thickness is 2mm. Then isothermal annealing is performed at 375℃ for 4h, and then solution treatment is performed at 493℃ for 11min, followed by immediate water quenching, quenching transfer time is 6s, and natural aging is performed for 255h.

[0178] The hot-rolled alloy is controlled by Brass texture, and the Brass texture strength of the hot-rolled alloy is 10.71, the Copper texture strength is 1.74, and the S texture strength is 1.74. The naturally aged alloy is controlled by Goss+Cube texture, and the Goss texture strength is 5.15, and the Cube texture strength is 5.15. Fig.32 and 33 Ph i 1, PHI, Constant Ph i2 represent the φ1, φ, φ2 angles in the Euler space. 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. Instability ΔK = 28.5MPa*m 1 / 2 Compared with Example 2, the Cu / Mg ratio is reduced and the Mg content is increased, which reduces the Goss texture strength in the aged alloy and increases the coarse second phase content in the matrix sharply. Not only is the ultimate tensile strength far from 500MPa, but the fatigue performance is also very poor. The fatigue crack growth rate in the TL direction is shown in Table 11. Fig.34 The vertical coordinate da / dN (mm / cyc le) represents the crack growth rate, and the horizontal coordinate DeltaK (MPa*m1 / 2) represents the stress intensity factor.

[0179] Table 11 Fatigue crack growth rate (×10 -4 mm / cycle)

[0180]

[0181] Combined with Figure 1-4It can be seen that after 470℃ / 48h+505℃ / 12h homogenization+95.6% reduction at 405℃, the hot-rolled texture is controlled by high-strength Copper, S and Brass textures, and their strengths are 8.20, 6.82 and 6.82 respectively, among which the Copper texture has the highest strength. After 340℃ / 5h annealing+485℃ / 30min solid solution+water quenching (transfer time 5s)+at least 240h natural aging, the hot-rolled plate is controlled by high-strength Goss and P textures, and the strengths of Goss and P textures are both 8.01, and the combined strength is 16.02. The expansion rate of the steady-state expansion zone of the alloy in Example 1 is relatively low, at ΔK=30MPa*m 1 / 2 When the expansion rate is only 15.2×10 -4 mm / cycle. The internal dislocation content of the matrix in the natural aging state is low, with only a small amount of coil dislocation.

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

[0183] from Fig. 9 It can be seen that after 465℃ / 36h+495℃ / 48h homogenization+435℃ 93% reduction+360℃ / 5h annealing+503℃ / 3min solid solution+water quenching (transfer time 3s)+at least 260h natural aging, the alloy of Example 3 is at ΔK=30MPa*m 1 / 2 When the expansion rate is 26.9×10 -4 mm / cycle.

[0184] Combined with Figure 10-14It can be seen that after 480℃ / 36h+495℃ / 24h homogenization+455℃ 85% reduction hot rolling+45% cold rolling, the texture is controlled by high-strength Brass texture, with Brass texture strength of 9.64, Copper texture strength of 3.18, and S texture strength of 3.18. After 365℃ / 6h annealing+500℃ / 10min solution treatment+water quenching (transfer time 7s)+at least 280h natural aging, the cold-rolled sheet is controlled by high-strength Goss texture, with Goss texture strength of 7.1, Cube texture strength of 4.1, and combined strength of 11.2. Example 3 alloy at ΔK=30MPa*m 1 / 2 When the expansion rate is 19.2×10 -4 mm / cycle. The tensile strength reaches 512.1MPa and the elongation is as high as 21.1%. The naturally aged microstructure contains Goss+Cube grain bands. The long diameter 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 naturally aged matrix is ​​composed of a small amount of coil dislocations + dislocation rings.

[0185] Combined with Figure 15-17 It can be seen that after 485℃ / 36h+495℃ / 36h homogenization+465℃ 91% reduction hot rolling+300℃ / 3h annealing+30% cold rolling texture+370℃ / 4h annealing+505℃ / 6min solid solution+water quenching (transfer time 10s)+at least 265h natural aging, the high-strength Goss+Cube texture is controlled, the Goss texture strength is 10.7, the Cube texture strength is 8.2, and the combined strength is 18.9. Example 4 alloy at ΔK=30MPa*m 1 / 2 When the expansion rate is 22.7×10 -4 mm / cycle. The interior of the natural aging matrix is ​​composed of coil dislocations + dislocation loops.

[0186] Combined with Figure 18-19 It can be seen that after 490℃ / 36h+490℃ / 36h homogenization+480℃ hot rolling with 91.5% reduction+415℃ / 3h annealing+495℃ / 8min solid solution+water quenching (transfer time 11s)+155℃ / 3h artificial aging. Example 5 alloy at ΔK=30MPa*m 1 / 2 When the expansion rate is 15.8×10 -4 mm / cycle. The interior of the aged matrix is ​​composed of high-density and moderate-sized Cu-Mg atomic clusters, and dislocations are basically eliminated.

[0187] Combined with Figure 20-21It can be seen that after 485℃ / 24h+490℃ / 48h homogenization+470℃ 87% reduction hot rolling+400℃ / 4h annealing+495℃ / 20min solid solution+water quenching (transfer time 6s)+173℃ / 1h artificial aging. Example 6 alloy at ΔK=30MPa*m 1 / 2 When the expansion rate is 29.1×10 -4 mm / cycle. The interior of the artificially aged matrix is ​​composed of large-sized and high-density Cu-Mg atomic clusters, and dislocations are basically eliminated.

[0188] Combined with Figure 22-25 It can be seen that after 475℃ / 36h+500℃ / 24h homogenization+490℃ 90% reduction, the hot rolling texture is controlled by high-strength Brass texture, with a strength of 12.23, and the strengths of Copper and S textures are both 1.99. After 390℃ / 4h annealing+495℃ / 15min solid solution+water quenching (transfer time 3s)+at least 260h natural aging, the hot rolled plate is controlled by strong Goss+Cube texture, with Goss texture strength of 7.87, Cube texture strength of 7.87, and Goss+Cube combined strength of 15.74. Comparative Example 1 alloy at ΔK=30MPa*m 1 / 2 When the expansion rate is 14.9×10 -4 mm / cycle. The naturally aged matrix is ​​composed of high-density dislocation loops.

[0189] Combined with Figure 26-29 It can be seen that after 475℃ / 34h+500℃ / 20h homogenization+486℃% reduction, the hot rolling texture is controlled by high-strength Brass texture, with Brass texture strength of 11.31, Copper texture strength of 5.63, and S texture strength of 3.73. After 385℃ / 4h annealing+495℃ / 15min solid solution+water quenching (transfer time 4s)+at least 260h natural aging, the hot rolled plate is controlled by strong Goss+Cube texture, with Goss texture strength of 9.96, Cube texture strength of 9.96, and Goss+Cube combined strength of 19.92. Comparative Example 2 alloy at ΔK=30MPa*m 1 / 2 The expansion rate is as high as 35×10 -4 mm / cycle. The naturally aged matrix is ​​composed of high-density dislocation loops.

[0190] Combined with Figure 30-31 It can be seen that after 490℃ / 35h+490℃ / 36h homogenization+480℃ hot rolling with 91% reduction+410℃ / 3h annealing+495℃ / 7min solid solution+water quenching (transfer time 10s)+159℃ / 3h artificial aging. 1 / 2The artificially aged matrix is ​​composed of high-density dislocation rings.

[0191] Combined with Figure 32-35 It can be seen that after 490℃ / 36h+495℃ / 24h homogenization+480℃ with 90% reduction, the hot-rolled texture is controlled by the high-strength Brass texture with a strength of 10.71, the Copper texture strength is 1.74, and the S texture strength is 1.74. After 375℃ / 4h annealing+493℃ / 11min solid solution+water quenching (transfer time 6s)+at least 255h natural aging, the hot-rolled plate is controlled by the Goss+Cube texture, with a Goss texture strength of only 5.15, a Cube texture strength of only 5.15, and a combined strength of 10.30. The expansion rate of the steady-state expansion zone of the alloy in Comparative Example 4 is relatively high, at ΔK=25MPa*m 1 / 2 The crack growth rate is as high as 71.6×10 - 4 mm / cycle. The naturally aged matrix is ​​composed of high density dislocation lines.

[0192] Combined with Figure 1-21 It can be seen that the tensile strength of the alloy of the present invention after homogenization, hot rolling or cold rolling after hot rolling, annealing, solution treatment and subsequent aging process treatment within the scope of the claims can be higher than 495MPa. Under such a high strength, at ΔK = 30MPa*m 1 / 2 The fatigue crack growth rate of the alloy is lower than 30×10 -4 mm / cycle. It exhibits excellent characteristics of high strength and fatigue resistance.

[0193] Combined with Figure 5-8 It can be seen from Figures 22-29 that when the Cu content in the alloy is less 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 high, the subsequent Goss and Cube texture strengths decrease. And the dislocation density in the natural aging state increases significantly. Although the fatigue performance is improved, the tensile strength is far from 495MPa. In addition, keeping the Cu and Mg content close and not adding Ag elements leads to the formation of high-density dislocation rings in the naturally aged alloy. There is no significant difference in the Goss and Cube texture strengths in the naturally aged state, but the tensile strength of the alloy is also lower than 495MPa, and the fatigue crack growth rate under high ΔK increases significantly.

[0194] Combined with Figure 10-17 It can be seen that before cold rolling, the hot-rolled plate was annealed at 300℃ / 3h. Compared with the unannealed plate before cold rolling, the Brass texture was enhanced, and then the Goss texture strength in the natural aging state was significantly enhanced, thereby increasing the instability ΔK.

[0195] Combined with Figure 18-19 It can be seen from 30-31 that the content of Cu and Mg elements is kept close and no Ag element is added. After under-aging, the alloy without Ag element is still composed of high-density dislocation rings, and its tensile strength is far from that of the alloy with Ag added, far below 495MPa. And the fatigue performance is also significantly reduced. The alloy without Ag addition has a ΔK = 26.6MPa*m 1 / 2 It has become unstable and fractured.

[0196] Combined with Figure 5-8 It can be seen from Figures 32-35 that when the Cu content in the alloy is not much different and the Mg content is higher than 1.4%, disordered high-density dislocations are formed in the naturally aged alloy, and the Goss and Cube texture strengths are significantly reduced. This not only significantly reduces the tensile strength of the naturally aged alloy, but also greatly deteriorates the fatigue properties, ΔK = 25MPa*m 1 / 2 The crack growth rate is as high as 71.6×10 -4 mm / cycle. When Cu and Mg elements are within the scope of the claims of the present invention, at ΔK = 25 MPa*m 1 / 2 When the crack growth rate is only 8.9×10 -4 mm / cycle.

[0197] By comparing the tensile mechanical properties and fatigue crack growth rates of Example 1-7 with those of Comparative Example 1-4, it can be seen that the alloy of Example 1-7 still has very excellent fatigue performance while ensuring high strength (>495MPa).

[0198] Finally, it should be noted that the terms "comprises," "includes," or any other variations thereof are intended to cover non-exclusive inclusion, such that a process, method, article, or apparatus that includes a series of elements includes not only those elements, but also other elements not explicitly listed, or also includes elements that are inherent to such process, method, article, or apparatus.

[0199] Although the preferred embodiments of the present invention have been described, those skilled in the art may make other changes and modifications to these embodiments once they have learned the basic creative concept. Therefore, the appended claims are intended to be interpreted as including the preferred embodiments and all changes and modifications that fall within the scope of the present invention.

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

Claims

1. A high-strength aluminum alloy with excellent fatigue damage resistance, characterized in that: In terms of mass fraction, the chemical composition of the high-strength aluminum alloy 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, 3≤Cu / Mg≤4; The microstructure of the high-strength aluminum alloy comprises Goss+Cube or Goss+P grain bands, the major diameter 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.

2. The high-strength aluminum alloy with excellent fatigue damage resistance according to claim 1, characterized in that: 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.

3. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 1 or 2, characterized in that: The preparation method comprises: An ingot is obtained by a smelting and casting process; The ingot is subjected to a primary homogenization treatment, and then the temperature is raised to perform a secondary homogenization treatment 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 is subjected to large deformation hot rolling, and then cold rolled or directly cold rolled after annealing to obtain a cold rolled sheet; Isothermally annealing the hot-rolled sheet or the cold-rolled sheet to obtain an annealed aluminum alloy sheet; 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; Wherein, the chemical composition of the ingot is as follows, measured by 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<0.06%, Si<0.06%, Zn<0.1%; the balance is Al, 3≤Cu / Mg≤4.

4. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 3, characterized in that: The ingot is subjected to a primary homogenization treatment, and then the temperature is raised to perform a secondary homogenization treatment to obtain a homogenized ingot, specifically comprising: The ingot is subjected to a primary homogenization treatment at 465-490° C. for 24-48 hours; After the primary homogenization treatment is completed, the temperature is raised to 490°C to 505°C for secondary homogenization treatment. The homogenization time is 12 to 48 hours to obtain a homogenized ingot.

5. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 3, characterized in that: The step of subjecting the homogenized ingot to large deformation hot rolling to obtain a hot-rolled plate specifically comprises: The homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm.

6. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 3, characterized in that: The homogenized ingot is subjected to large deformation hot rolling, and then cold rolled or directly cold rolled after annealing to obtain a cold rolled sheet, specifically comprising: The homogenized ingot is subjected to large deformation hot rolling at 390-490° C., with a total rolling deformation greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm, and then the hot-rolled plate is directly cold-rolled 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 greater than 85%, to obtain a hot-rolled plate with a thickness of 2-6 mm, the hot-rolled plate is annealed at a temperature of 260-330° C. for 2-4 h, and the annealed plate is subsequently cold rolled with a cold rolling deformation of ≤50% to obtain a cold-rolled plate.

7. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 3, characterized in that: The isothermal annealing of the hot-rolled sheet or the cold-rolled sheet to obtain the annealed aluminum alloy sheet specifically includes: Isothermally annealing the hot-rolled sheet or the cold-rolled sheet to obtain an annealed aluminum alloy sheet; Among them, when the large deformation hot rolling temperature is lower than 450°C, the isothermal annealing temperature is controlled at 320°C ~ 370°C, and the annealing time is 2 to 5 hours. When the large deformation hot rolling temperature is higher than 450°C, the isothermal annealing temperature is controlled at 340°C ~ 420°C, and the annealing time is 3 to 6 hours.

8. The method for preparing a high-strength aluminum alloy with excellent fatigue damage resistance according to claim 3, characterized in that: 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, specifically comprising: The annealed aluminum alloy plate is subjected to high temperature short-time solution treatment, the solution treatment temperature is 475-505° C., the time is 3-30 minutes, and then water quenching is performed with water as the medium, the quenching transfer time is ≤15 seconds, and then aging treatment is performed to obtain a high-strength aluminum alloy; Wherein, when the aging treatment adopts the natural aging system, the natural aging time is ≥240h, and when the artificial aging system is adopted, the artificial aging temperature is 140°C to 180°C, and the aging time is 1 to 5h.

9. Use of a high-strength aluminum alloy with excellent fatigue damage resistance as claimed in claim 1 or 2 in the preparation of aviation or aerospace vehicle structural parts.

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