High-strength alloy material and preparation process thereof

Abstract

The application relates to the technical field of alloy materials, in particular to a high-strength alloy material and a preparation process thereof. The application overcomes the problem that the mechanical properties and high-temperature resistance of alloy materials in the prior art cannot be considered simultaneously. The alloy component obtained by the application comprises Al-Li-Cu-Zr-Sc-Yb-In-Ag, and a hot-pressed alloy is obtained through high-temperature zirconium melting-low-temperature lithium pressing smelting, hierarchical homogenization, deformation heat processing, deep cooling and pre-strain dislocation network induction technology; then a bimodal structure coexisting with equiaxed fine crystals and elongated fiber crystals is constructed through gradient instantaneous annealing, and a dense carbon atom layer is deposited on the surface by combining chemical vapor deposition technology. The alloy material obtained by the application has good fatigue resistance and corrosion resistance, and is suitable for large thin-walled stressed components of aerospace vehicles and high-temperature protection fields.

Description

Technical Field

[0001] This invention relates to the field of alloy materials technology, specifically to a high-strength alloy material and its preparation process. Background Technology

[0002] Aluminum-lithium alloys are widely recognized as core structural materials for achieving lightweight and high-performance in next-generation aerospace equipment due to their low density, high elastic modulus, and excellent cryogenic performance. These include hypersonic vehicle skins, cryogenic fuel tanks for launch vehicles, and large load-bearing components. However, under extreme service conditions, these materials not only need to withstand enormous mechanical loads but also endure continuous thermal stress caused by aerodynamic heating and chemical corrosion in harsh service environments.

[0003] Existing aluminum-lithium alloy technologies still face unresolved issues in practical applications: traditional strengthening mechanisms, in pursuit of ultra-high strength, often lead to a sharp decline in fracture toughness and plasticity. This imbalance in strength-plasticity product makes the material highly susceptible to brittle fracture under impact loads or complex stress states. In addition, in high-temperature environments exceeding 150°C, the metastable strengthening phases in existing alloys rapidly undergo over-aging coarsening and tend to dissolve or segregate at grain boundaries. This microstructural instability leads to severe thermal softening, causing a precipitous drop in the tensile strength at high temperatures and the strength retention rate after thermal exposure, making it difficult to meet the safety requirements of hypersonic service environments.

[0004] Meanwhile, the strong texture produced by traditional processes leads to severe mechanical anisotropy, limiting the structural design of complex irregular parts. The extremely high chemical reactivity of aluminum-lithium alloys makes them highly susceptible to salt spray environments, resulting in pitting corrosion, exfoliation corrosion, and stress corrosion cracking. Existing surface protection methods often fail to provide a dense and highly bonded physical shielding layer while maintaining the stability of the alloy's reinforced microstructure. In summary, current technologies continuously improve the overall performance of alloy materials through component innovation and process modification; however, the challenge of simultaneously achieving both mechanical properties and high-temperature resistance remains, becoming a core requirement in the current high-end equipment manufacturing field.

[0005] To address this, a high-strength alloy material and its preparation process are proposed. Summary of the Invention

[0006] The purpose of this invention is to provide a high-strength alloy material and its preparation process. The alloy composition obtained by this invention comprises Al-Li-Cu-Zr-Sc-Yb-In-Ag. It is hot-pressed as an alloy through high-temperature zirconium melting-low-temperature lithium pressing, graded homogenization, deformation hot working, cryogenic treatment, and pre-strain induced dislocation network processes. Then, a bimodal microstructure consisting of equiaxed fine grains and elongated fibrous grains is constructed through gradient instantaneous annealing, and a dense carbon atom layer is deposited on the surface using chemical vapor deposition. The alloy material obtained by this invention possesses excellent fatigue resistance and corrosion resistance, making it suitable for large, thin-walled load-bearing components in aerospace vehicles and for high-temperature protection applications.

[0007] To achieve the above objectives, the present invention provides the following technical solution:

[0008] This invention provides a preparation process for a high-strength alloy material, comprising the following preparation steps:

[0009] High-purity aluminum ingots are placed in a vacuum induction melting furnace, and a vacuum of 10°C is applied. -2 After Pa, high-purity argon gas is introduced for protection, and the temperature is raised to 775℃-790℃ to completely melt the aluminum ingot. Then, aluminum-zirconium master alloy, aluminum-scandium master alloy, and aluminum-ytterbium master alloy are added sequentially according to mass percentage. The electromagnetic stirring system is started with a power frequency of 10Hz and stirring is continued for 15-20 minutes to obtain a homogeneous elemental liquid. Low-frequency stirring can provide a greater penetration depth, ensuring a strong macroscopic circulation throughout the molten pool and ensuring that the high-melting-point microalloying elements achieve atomic-level uniform dispersion in the liquid phase. The temperature of the homogeneous elemental liquid is reduced to 730℃-745℃ through a circulating cooling system, and elemental copper particles, silver particles, and indium sheets are added. The power frequency is maintained and stirring is carried out for 10 minutes. The thermodynamic energy within this temperature range promotes the initial construction of coherent precursor clusters of solute atoms. Next, the furnace temperature is lowered to 695℃-710℃. A special bell-shaped feeder is used to pre-fill the bell-shaped container with 0.04MPa of high-purity argon gas to form a positive pressure buffer layer. The aluminum-lithium master alloy is steadily pressed into the liquid at a speed of 200mm / s to a depth of 300mm-500mm below the liquid surface. The air cushion pressure is used to suppress molten splashing and isolate the air, avoiding severe oxidation and burning of lithium elements. After the melt is completely stable, high-purity argon gas is introduced for rotary powder spraying refining. The flow rate is controlled at 15-25L / min, the rotor speed is 500rpm, and the refining time is 20-25min. After standing for 15min, semi-continuous casting is carried out. The casting speed is controlled at 40-60mm / min, and the circulating water temperature is 15℃-25℃ to obtain a uniform alloy ingot.

[0010] The alloy ingot is placed in a segmented heating furnace for graded homogenization to obtain a treated alloy. In the first stage, the alloy is held at 455-465℃ for 12-14 hours to eliminate low-melting-point eutectic segregation. In the second stage, the temperature is raised to 515-535℃ and held for 24-28 hours to allow Sc and Zr atoms to disperse and form dispersed nanoscale Al3(Sc,Zr,Yb) composite phase particles. Subsequently, the treated alloy is subjected to deformation hot working. The billet temperature is set at 425℃-440℃, and the hot-pressed alloy is obtained through hot extrusion. The single-pass reduction rate is controlled at 15%-25%, the total deformation reaches 85%-92%, and the final processing temperature is strictly limited to above 365℃ to retain high-density dislocations and deformation substructures.

[0011] The hot-pressed alloy was solution-treated at 530-540℃ for 1.5 hours; then rapidly immersed in circulating cold water, with the quenching transfer time controlled within 5 seconds to ensure the integrity of the supersaturated solid solution; after quenching, the hot-pressed alloy was placed in an ultra-low temperature cryogenic chamber containing liquid nitrogen for no more than 30 minutes, and subjected to extreme cryogenic treatment at -196℃ for 3-4 hours, utilizing the huge thermal stress gradient to increase the matrix vacancy concentration and lattice distortion energy; after returning to room temperature, within 1 hour, a 5% permanent plastic pre-strain was applied using a CNC stretching machine to obtain a stretched alloy, artificially introducing a high-density, regularly distributed dislocation network; the stretched alloy was then subjected to a high-power induction heating device. The material is subjected to gradient annealing with a heating rate set at 18℃ / s-25℃ / s. It is rapidly heated to 485-490℃ and held for 60-80s before being immediately subjected to a second water quench to obtain the annealed alloy. A high-frequency induction heating device or a multi-segment infrared rapid heating furnace is used, along with an infrared multi-point temperature measurement system for real-time feedback, to ensure that the heating rate of the surface and core of large components is synchronized. The annealed alloy is subjected to a three-stage aging treatment to obtain the alloy precursor. The first stage involves pre-nucleation at 110-120℃ for 8 hours, the second stage involves heating to 135-145℃ for 10 hours of phase growth, and the third stage involves heating to 160-170℃ for 12 hours of phase stabilization.

[0012] The alloy precursor was ultrasonically cleaned with anhydrous ethanol for 15 minutes, dried, and then placed into a chemical vapor deposition reaction chamber. The ultrasonic power was 400 W and the ultrasonic frequency was 30 kHz. The reaction chamber was then evacuated to a vacuum of 8.0 × 10⁻⁶. -4Below Pa, argon gas is introduced at a flow rate of 50-80 sccm for vacuum pretreatment; the RF power supply is set to 150-250W, bias voltage -300V, and in-situ plasma sputtering cleaning of the alloy surface is performed for 10-15 minutes using argon plasma to obtain the pretreated alloy; residual oxide film on the surface is removed to improve the adhesion of the subsequent carbon layer; acetylene and hydrogen are introduced into the pretreated alloy, with the acetylene flow rate at 15-30 sccm (standard milliliters per minute) and the hydrogen flow rate... The deposition rate is 5-10 sccm, the working pressure is controlled at 5-15 Pa, the radio frequency power is 300-500 W, the deposition time is 30-60 min, and the deposition temperature is 150℃-170℃, resulting in a carbon atom deposition layer with a thickness of 200-500 nm. The alloy material is obtained by cooling it to below 50℃ under argon protection. EBSD characterization shows that the equiaxed fine grain size inside the alloy is 0.5-2 μm, the length-to-diameter ratio of the elongated fiber is greater than 10:1, and the two are alternately distributed.

[0013] In this invention, acetylene and hydrogen are introduced into the pretreated alloy, and non-equilibrium plasma generated by a radio frequency power supply is used to deeply dissociate acetylene molecules, enabling carbon atoms to achieve dense deposition in a non-thermal equilibrium state at a low temperature of 150℃-170℃. The resulting carbon atom deposition layer has a diamond-like structure and is constructed simultaneously with the aging process at 150℃-170℃. The carbon atoms fill the microscopic defects on the surface and form chemical bonds, ensuring the density and bonding strength of the surface physical shielding layer.

[0014] The present invention also provides a high-strength alloy material, which, by mass percentage, comprises 2.2-2.6% lithium, 3.2-3.8% copper, 0.12-0.18% scandium, 0.05-0.09% indium, 0.35-0.45% silver, 0.1-0.15% zirconium, 0.01-0.05% ytterbium, with the balance being aluminum.

[0015] Compared with the prior art, the beneficial effects of the present invention are as follows:

[0016] 1. This invention introduces microalloying elements such as Sc, Zr, and Yb and combines them with graded homogenization treatment to form highly thermally stable composite nanoparticles in the matrix. These nanoparticles can maintain a coherent relationship with the aluminum matrix and have extremely high coarsening resistance. They can still firmly pin dislocations and subgrain boundaries in high-temperature environments, inhibiting the recovery and recrystallization of the matrix, thus meeting the microstructure stability requirements in hypersonic service environments.

[0017] 2. This invention successfully constructs a heterogeneous structure within the alloy by employing a gradient instantaneous annealing process, where submicron equiaxed fine grains and micron-scale elongated fibrous grains coexist. The equiaxed fine grain region can effectively store dislocations and provide excellent plastic deformation capabilities, while the fibrous grain region, which retains high-density dislocations and deformation substructures, provides extremely high matrix strength. Combined with high-density, dispersed nanoscale reinforcing phases induced by elements such as In and Ag, the synergistic improvement of alloy strength and plasticity is achieved, solving the problem of insufficient toughness in traditional alloys.

[0018] 3. This invention introduces a heterogeneous bimodal structure to create multiple obstacles for crack propagation. When fatigue cracks pass through the heterogeneous interface between equiaxed and fibrous crystals, frequent crack deflection and passivation occur, significantly increasing the crack propagation path and reducing its propagation rate. Combined with a high-modulus carbon atom deposition layer on the surface, it effectively constrains the slip of surface dislocations, inhibits the early initiation of fatigue cracks on the surface, greatly improves the fatigue limit of the alloy, and extends the fatigue service life of the structural components.

[0019] 4. This invention introduces a chemical vapor deposition layer of carbon atoms to form a dense and chemically inert nanoscale physical barrier on the alloy surface. This carbon layer not only significantly improves the surface's resistance to exfoliation corrosion and pitting corrosion, but its extremely high density also plays an excellent role in oxygen barrier, effectively preventing oxygen atoms in the environment from penetrating into the matrix and causing internal oxidation and surface burn-off of lithium elements, thereby enhancing the alloy's self-passivation ability and chemical stability in complex and harsh environments.

[0020] 5. This invention solves the problems of lithium element burn-off and high-melting-point element segregation by three-stage temperature-controlled melting, and controls the carbon deposition temperature to match the peak aging temperature of the alloy, so as to achieve simultaneous surface modification and microstructure stabilization. Through the deep coupling of multiphysics field and chemical vapor deposition, it not only ensures that the matrix strengthening phase does not undergo over-aging softening, but also further improves its application in aerospace equipment from composition design, atomic cluster regulation, grain heterogeneity to surface carbon layer protection. Attached Figure Description

[0021] Figure 1 The graph shows the change in the heat exposure tensile strength retention rate of Examples 1-5 and Comparative Examples 1-5 of the present invention. Detailed Implementation

[0022] The technical solutions of the embodiments of the present invention will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0023] In this invention, the high-purity aluminum ingot has a purity greater than 99.99%, a melting point of 660.3℃, and a density of 2.70 g / cm³; the lithium content in the aluminum-lithium master alloy is 10% by mass, and the moisture content is less than 100 ppm; the copper particles have a particle size of 3-8 mm; the silver particles have a particle size of 1-3 mm; the zirconium content in the aluminum-zirconium master alloy is 5% by mass; the scandium content in the aluminum-scandium master alloy is 2% by mass; and the ytterbium content in the aluminum-ytterbium master alloy is 1% by mass.

[0024] Please see Figure 1 This invention provides a high-strength alloy material and its preparation process, the technical solution of which is as follows:

[0025] Example 1

[0026] High-purity aluminum ingots are placed in a vacuum induction melting furnace, and a vacuum of 10°C is applied. -2 After Pa, high-purity argon gas is introduced for protection, and the temperature is raised to 780℃ to completely melt the aluminum ingot. Then, aluminum-zirconium master alloy, aluminum-scandium master alloy, and aluminum-ytterbium master alloy are added sequentially according to mass percentage. The electromagnetic stirring system is started with a power frequency of 10Hz and stirring is continued for 20 minutes to obtain a homogeneous elemental liquid. The temperature of the homogeneous elemental liquid is reduced to 740℃ through a circulating cooling system, and elemental copper particles, silver particles, and indium metal sheets are added. The power frequency is maintained, and stirring is continued for 10 minutes. The furnace temperature is then reduced to 700℃, and aluminum-lithium master alloy is steadily pressed into the liquid surface at a speed of 200mm / s using a special bell jar feeder. At a distance of 400mm; after the melt has completely stabilized, high-purity argon gas is introduced for rotary powder spraying refining, with the flow rate controlled at 20L / min, the rotor speed at 500rpm, and the refining time at 25min. After standing for 15min, semi-continuous casting is carried out, with the casting speed controlled at 50mm / min and the circulating water temperature at 20℃, to obtain an alloy ingot with uniform composition; calculated by mass percentage, it contains 2.4% lithium, 3.5% copper, 0.15% scandium, 0.07% indium, 0.4% silver, 0.12% zirconium, 0.03% ytterbium, and the balance being aluminum.

[0027] Alloy ingots were placed in a segmented heating furnace for graded homogenization to obtain treated alloys. In the first stage, the alloys were held at 460℃ for 13 hours to eliminate low-melting-point eutectic segregation. In the second stage, the temperature was raised to 525℃ and held for 26 hours. Subsequently, the treated alloys were subjected to deformation hot working. The billet temperature was set at 430℃, and the hot-pressed alloys were obtained through hot extrusion. The single-pass reduction rate was controlled at 20%, the total deformation reached 88%, and the final processing temperature was strictly limited to above 365℃ to retain high-density dislocations and deformation substructures.

[0028] The hot-pressed alloy was solution-treated at 530℃ for 1.5 hours; then rapidly immersed in circulating cold water, with the quenching transfer time controlled within 5 seconds to ensure the integrity of the supersaturated solid solution; after quenching, the hot-pressed alloy was placed in an ultra-low temperature cryogenic chamber containing liquid nitrogen for 30 minutes for extreme cryogenic treatment at -196℃ for 3 hours; after returning to room temperature, within 1 hour, a 5% permanent plastic pre-strain was applied using a CNC stretching machine to obtain a stretched alloy; the stretched alloy was then subjected to gradient annealing using a high-power induction heating device, with a heating rate set at 18℃ / s, rapidly heated to 488℃ and held for 70 seconds, followed immediately by a second water quench to obtain an annealed alloy; the annealed alloy underwent a three-stage aging treatment to obtain the alloy precursor: the first stage involved pre-nucleation at 110℃ for 8 hours, the second stage involved phase growth at 140℃ for 10 hours, and the third stage involved stabilizing the phase at 168℃ for 12 hours.

[0029] The alloy precursor was ultrasonically cleaned with anhydrous ethanol for 15 minutes, dried, and then placed into a chemical vapor deposition reaction chamber. The ultrasonic power was 400 W and the ultrasonic frequency was 30 kHz. The reaction chamber was then evacuated to a vacuum of 8.0 × 10⁻⁶. -4 The pressure was below Pa, and then argon gas was introduced at a flow rate of 60 sccm for vacuum pretreatment. The RF power was set to 200W and the bias voltage to -300V. The alloy surface was then subjected to in-situ plasma sputtering cleaning for 15 minutes using argon plasma to obtain the pretreated alloy. The residual oxide film on the surface was removed to improve the adhesion of the subsequent carbon layer. Acetylene was introduced into the pretreated alloy at a flow rate of 22 sccm and hydrogen at a flow rate of 8 sccm. The working pressure was controlled at 10 Pa, the RF power was 400W, the deposition time was 50 minutes, and the temperature was 160℃ to obtain a carbon atom deposition layer with a thickness of 350 nm. The alloy material was obtained by cooling it to below 50℃ in the furnace under argon protection.

[0030] Examples 2-5 follow the same preparation method and parameter conditions as Example 1, with differences shown in Table 1.

[0031]

[0032] Comparative Example 1 is the same as Example 1, except that trace elements Sc, Zr and Yb are not added during the preparation of the alloy ingot, while the amounts of other components remain unchanged.

[0033] Comparative Example 2 is the same as Example 1, except that trace elements In and Ag are not added during the preparation of the alloy ingot, while the amounts of other components remain unchanged.

[0034] Comparative Example 3 is the same as Example 1, except that no trace element Yb is added during the preparation of the alloy ingot, while the amounts of other components remain unchanged.

[0035] Comparative Example 4 is the same as Example 1, except that no trace element Sc is added during the preparation of the alloy ingot, while the amounts of other components remain unchanged.

[0036] Comparative Example 5 is the same as Example 1, except that no trace element In is added during the preparation of the alloy ingot, while the amounts of other components remain unchanged.

[0037] Comparative Example 6 is the same as Example 1, except that the alloy material is not subjected to -196℃ cryogenic treatment during preparation, and is directly pre-stretched after solution quenching.

[0038] Comparative Example 7 is the same as Example 1, except that the melting and stirring speed is reduced to 100 rpm during the alloy ingot preparation process.

[0039] Comparative Example 8 is the same as Example 1, except that the total deformation during hot working is adjusted to 30%.

[0040] Comparative Example 9 is the same as Example 1, except that conventional homogenization annealing at 450°C for 2 hours is used instead of gradient instantaneous annealing process, while the rest of the process remains unchanged.

[0041] Comparative Example 10 is the same as Example 1, except that the annealing heating rate is reduced to 2°C / s, while the rest of the process remains unchanged.

[0042] Comparative Example 11 is the same as Example 1, except that the alloy precursor is not treated by chemical vapor deposition process, and the alloy material is obtained after alloy aging treatment.

[0043] Comparative Example 12 is the same as Example 1, except that the aging treatment uses a single 165°C aging treatment and does not include the pre-nucleation and phase growth stages.

[0044] Comparative Example 13 is the same as Example 1, except that the thickness of the carbon layer deposited by chemical vapor deposition is 100 nm.

[0045] Experiment Example 1: Thermal Stability Test

[0046] The alloy materials prepared in Examples 1-5 and Comparative Examples 1-5 were subjected to thermal stability tests. The tensile strength of the alloy materials was tested according to GB / T228.1-2021, and the high-temperature tensile properties were tested according to GB / T 33886-2017. The alloy materials were heated to 200℃ and held for 10 minutes before tensile testing to obtain the high-temperature tensile strength. The alloys were also exposed to 200℃ for 100 hours, cooled to room temperature, and then the tensile strength was tested according to GB / T 228.2-2015. The strength retention rate was calculated by the ratio of the tensile strength before and after the exposure. The test results are shown in Table 2. The changes in the thermal exposure tensile strength retention rate of Examples 1-5 and Comparative Examples 1-5 are as follows: Figure 1 As shown.

[0047]

[0048] Through Table 2, Figure 1 The results show that the thermal stability of the alloy material obtained in the comparative example, through adjustments to the composition and process, is significantly different from that in the examples. In the examples, the introduction of microalloying elements such as Sc, Zr, and Yb allows these elements to interact during the hierarchical homogenization and subsequent heat treatment processes, forming dispersed nanoscale Al3(Sc,Zr,Yb) composite phase particles. These particles can maintain a coherent or semi-coherent relationship with the matrix, exhibiting extremely high thermal stability. They effectively pin dislocations and subgrain boundaries, suppressing the recovery and recrystallization process of the alloy at 200℃. Additionally, the introduction of trace amounts of... Elements In and Ag have strong binding energy with vacancies in the matrix, which can effectively promote the uniform precipitation of Al2CuLi strengthening phase in the crystal. This synergistic effect of multi-scale strengthening phases enables the examples to maintain extremely high tensile strength and strength retention rate at high temperatures. In Comparative Example 1, since no trace elements Sc, Zr and Yb are added during the preparation process, the alloy lacks the key heat-resistant strengthening phase. At a high temperature of 200°C, without the pinning of grain boundaries by composite phase particles, the alloy is prone to grain coarsening and recovery, resulting in a significant decrease in its high-temperature strength.

[0049] Based on the results of Comparative Examples 2-5, it can be seen that in the absence of trace elements In and Ag, the alloy loses the key mechanisms for trapping vacancies and promoting precipitation, resulting in insufficient nucleation kinetics of the main strengthening phase. The precipitated phase tends to coarsen at grain boundaries and cannot form effective dislocation barriers within the grains. During heating, the coarse grain boundary precipitates are prone to dissolution or aggregation, severely affecting the thermal exposure stability of the alloy. Yb not only has the function of purifying grain boundaries, but can also form more stable composite core-shell structure particles with Sc and Zr. The lack of Yb slightly reduces the coarsening resistance of nanoparticles, leading to a decrease in the alloy's long-term stability. The strength retention rate after heat exposure is reduced; although Zr can form Al3Zr particles after the absence of Sc, their density and pinning strength to dislocations are far inferior to those of the composite phase particles, which means that the main anti-recrystallization driving force is missing, resulting in the alloy's tensile strength at 200℃ being significantly lower than that of all embodiments; the addition of In atoms can significantly improve the dispersion of precipitated phases in the early stage of aging, but the absence of In hinders the formation of vacancy-solute clusters, resulting in uneven phase distribution and easy coarsening, which greatly reduces its strengthening effect at high temperatures. Although there are still Sc / Zr particles to support it, the overall thermal stability is reduced.

[0050] Experiment Example 2: Mechanical Property Testing

[0051] The alloy materials prepared in Examples 1-5 and Comparative Examples 5-9 were subjected to mechanical property tests. In accordance with GB / T228.1-2021, standard tensile specimens in the L direction (longitudinal) and LT direction (transverse) were processed using an electronic universal tensile testing machine. The tensile strength, specified plastic extension strength and elongation at break were determined by controlling the strain rate. The test results are shown in Table 3.

[0052]

[0053] As shown in Table 3, the mechanical properties of the alloy materials obtained in the comparative examples, through adjustments to the composition and process, are significantly different from those in the examples. The introduction of trace elements In and Ag in the examples allows them to form atomic clusters by combining with vacancies. These clusters act as heterogeneous nucleation centers during aging precipitation, resulting in extremely dispersed precipitates and significantly improving the tensile and yield strength of the matrix. Furthermore, the composite nanoparticles formed by Sc, Zr, and Yb can pin dislocations and grain boundaries during hot working and annealing, inhibiting complete recrystallization of the matrix and preserving high-density dislocations and deformation substructures, which is beneficial for maintaining extremely high strength. The gradient instantaneous annealing process constructs a bimodal heterogeneous structure with both equiaxed fine grains and elongated fibrous grains, achieving a synergistic improvement in strength and plasticity.

[0054] In Comparative Example 5, the lack of In, a key element for inducing precipitation, significantly reduced the nucleation kinetics of the strengthening phase. The precipitated phase tended to coarsen at grain boundaries, resulting in insufficient intragranular strengthening and a significant decrease in plasticity. In Comparative Example 6, without cryogenic treatment at -196℃, the matrix could not generate a high concentration of lattice distortion and vacancies through a large thermal stress gradient. This weakened the driving force for precipitated phases during subsequent aging, reduced precipitation density, and consequently weakened yield strength. In Comparative Example 7, when the stirring speed was reduced to 100 rpm, the high-melting-point microalloying elements could not achieve atomic-level uniform dispersion in the liquid phase, resulting in significant micro-compositional segregation, which hindered subsequent aging. The uneven distribution of precipitates makes it prone to stress concentration and premature fracture during stretching. In Comparative Example 8, when the total deformation during hot working is adjusted to 30%, the strain energy stored in the matrix is ​​low, and it is impossible to retain a high-density dislocation network. Furthermore, it is difficult to form an ideal bimodal heterostructure during the subsequent instantaneous annealing process, resulting in a significant reduction in the strengthening increment of the material. Although the elongation has recovered somewhat, the strength is far lower than that of the example. In Comparative Example 9, conventional annealing is used instead of gradient instantaneous annealing. Due to the long holding time, the matrix undergoes complete recrystallization, losing the support for strength provided by the bimodal structure composed of elongated fibrous crystals and fine equiaxed crystals, leading to a complete collapse of mechanical properties.

[0055] Experiment Example 3: Fatigue Resistance Test

[0056] The alloy materials prepared in Examples 1-5 and Comparative Examples 7-11 were subjected to fatigue resistance tests according to GB / T3075-2021, on a high-frequency fatigue testing machine, with a stress ratio R=0.1 and a frequency of 100Hz. The tests were conducted over 10 hours. 7 The fatigue limit under cycle number was determined; standard CT specimens were used to observe the deflection behavior of cracks passing through the equiaxed / fiber-like heterostructure interface; the test results are shown in Table 4.

[0057]

[0058] As shown in Table 4, the fatigue resistance of the alloy material obtained in the comparative example, through adjustments to the composition and process, is significantly different from that in the example. In the example, by introducing trace elements Sc, Zr, and Yb and implementing a gradient instantaneous annealing process, a unique bimodal heterogeneous structure with equiaxed fine grains and elongated fibrous grains was constructed inside the alloy. This structure can effectively induce crack deflection, increase the crack propagation path, and significantly reduce the fatigue crack propagation rate. In addition, the carbon atom layer deposited on the surface acts as a dense physical barrier, effectively constraining the slip of surface dislocations, inhibiting the premature initiation of fatigue cracks on the surface, and greatly improving the fatigue limit.

[0059] In Comparative Example 7, when the melting stirring speed was reduced to 100 rpm, high-melting-point microalloying elements such as Sc and Zr could not achieve atomic-level dispersion in the aluminum melt, resulting in micro-segregation and the formation of coarse primary intermetallic compound particles in the ingot. These coarse particles easily become stress concentration points under cyclic loading, inducing crack initiation and leading to a significant decrease in fatigue limit. In Comparative Example 8, when the total deformation during hot working was only 30%, the dislocation density and strain energy accumulated inside the material were insufficient. Due to the lack of deflection and passivation of cracks by fine-grained regions, the crack propagation in a single microstructure was less resistant, resulting in a faster crack propagation rate. In Comparative Example 9, conventional homogenization annealing was used, causing complete recrystallization of the alloy and the disappearance of the original elongated fibrous structure. Due to the loss of this soft / hard phase... The heterogeneous structure effectively controls the crack propagation path, significantly reducing the material's damage tolerance, fatigue limit, and crack propagation resistance. In Comparative Example 10, when the annealing heating rate is reduced to 2℃ / s, the slow heating process allows sufficient time for grains to grow and merge before reaching the target temperature. Grain coarsening significantly weakens the grain boundaries' ability to block dislocation slip, not only reducing the material's strength but also making fatigue cracks more likely to initiate and propagate within the coarse grains. In Comparative Example 11, although the excellent microstructure of the matrix is ​​maintained, the lack of a carbon atom layer to fill surface micro-defects and anchor dislocation overflow makes the alloy surface more susceptible to the synergistic damage of environmental oxidation and cyclic stress, leading to a decrease in the fatigue limit. This also demonstrates the key contribution of the surface modification layer to improving long-life fatigue performance.

[0060] Experimental Example 4: Corrosion Resistance

[0061] The alloy materials prepared in Examples 1-5, Comparative Examples 4-5, and Comparative Examples 11-13 were subjected to corrosion resistance tests. Following GB / T 15970.6-2023 and GB / T 22639-2022, slow strain rate tensile testing was used, and the stress corrosion cracking threshold was determined in a 3.5% NaCl solution. The alloy samples were immersed in an acidic salt solution, and the surface delamination was observed and rated periodically. EA indicates slight delamination, EB indicates moderate delamination, EC indicates severe delamination, and ED indicates extremely severe delamination. The test results are shown in Table 5.

[0062]

[0063] As shown in Table 5, the corrosion resistance of the alloy material obtained in the comparative example, through adjustments to the composition and process, is significantly different from that in the example. In the example, the composite phase nanoparticles not only act as grain boundary pinning agents but also reduce the driving force of electrochemical corrosion by adjusting the grain boundary energy level. In addition, the introduced trace elements In and Ag form stable atomic clusters with vacancies in the alloy, serving as non-uniform nucleation sites and inducing their dispersion in the grain as extremely fine plates, effectively breaking the formation of continuous precipitation films at the grain boundaries and blocking the anodic channel for stress corrosion crack propagation along the grain boundaries. The dense carbon atom layer deposited on the surface endows the alloy material with excellent physical shielding properties, and the extremely high chemical stability of the carbon layer significantly enhances the surface's self-passivation ability.

[0064] In Comparative Example 4, without the crucial refining and pinning component Sc, the alloy cannot form sufficiently dense composite phase nanoparticles during homogenization, resulting in coarse grains and high grain boundary energy. In corrosive media, unprotected grain boundaries are prone to preferential oxidation and exfoliation, leading to a worse exfoliation corrosion rating. In Comparative Example 5, when the trace element In, a solute scavenger, is missing from the composition, uniform induced precipitation cannot be achieved within the matrix through vacancy clustering. This results in a large enrichment of Cu and Li atoms, forming coarse, continuous precipitated films at grain boundaries, creating a significant potential difference with the adjacent matrix, causing severe stress corrosion sensitivity and a significant reduction in the stress corrosion cracking threshold. In Comparative Example 11, due to the lack of chemical vapor deposition (CVD) treatment, the alloy surface is directly exposed to the external salt spray environment. Although the matrix... While the structure possesses some resistance, in the absence of a physical barrier of a high-modulus, chemically inert carbon layer, chloride ions can rapidly penetrate the surface oxide film and attack subsurface defects, leading to a significant increase in the depth of exfoliation corrosion. Combined with the results of Comparative Example 13, it is evident that due to the excessively thin carbon layer, discontinuous pinhole defects are easily generated during film formation. In stress corrosion testing, these micron-sized pinholes become microchannels for the electrolyte solution to penetrate the substrate, and the resulting small anode-large cathode effect accelerates localized corrosion penetration, reducing corrosion resistance. In Comparative Example 12, the use of a single aging treatment without a segmented pre-nucleation process led to uncontrolled precipitate growth. The single-stage high-temperature aging resulted in severe over-aging coarsening at grain boundaries, increasing the width of the copper-poor region at the grain boundaries and forming electrochemically weak zones that are highly susceptible to corrosion, significantly reducing corrosion resistance.

[0065] Although embodiments of the invention have been shown and described, it will be understood by those skilled in the art that various changes, modifications, substitutions and alterations can be made to these embodiments without departing from the principles and spirit of the invention, the scope of which is defined by the appended claims and their equivalents.

Claims

1. A preparation process for a high-strength alloy material, characterized in that, The preparation process includes the following steps: aluminum ingots, aluminum-zirconium master alloys, aluminum-scandium master alloys, aluminum-ytterbium master alloys, copper particles, silver particles, indium metal sheets, and aluminum-lithium master alloys are melted in stages and cast to obtain alloy ingots; the alloy ingots are subjected to graded homogenization treatment and then deformed hot working to obtain hot-pressed alloys; the hot-pressed alloys are subjected to solution treatment, quenching followed by cold treatment, plastic pre-straining and gradient annealing, and then subjected to three-stage aging treatment to obtain alloy precursors; the alloy precursors are subjected to chemical vapor deposition to obtain the alloy material; wherein the chemical vapor deposition uses a mixed gas of acetylene and hydrogen for carbon layer deposition.

2. The preparation process of a high-strength alloy material according to claim 1, characterized in that, The preparation of the alloy ingot includes the following steps: placing a high-purity aluminum ingot in a vacuum induction melting furnace, evacuating and filling it with argon gas for protection, so that the aluminum ingot is completely melted; then adding the aluminum-zirconium master alloy, the aluminum-scandium master alloy, and the aluminum-ytterbium master alloy in sequence according to mass percentage, and stirring to obtain a homogeneous elemental liquid; adding the elemental copper particles, the silver particles, and the indium sheet to the homogeneous elemental liquid, and continuing to stir; then cooling down and pressing the aluminum-lithium master alloy below the liquid surface; introducing high-purity argon gas for rotary powder spraying refining, and obtaining the alloy ingot with uniform composition through semi-continuous casting.

3. The preparation process of a high-strength alloy material according to claim 1, characterized in that, The preparation of the hot-pressed alloy includes the following steps: placing the alloy ingot in a segmented heating furnace for graded homogenization treatment to obtain the treated alloy; the first stage temperature of the graded homogenization treatment is 455-465℃; the second stage temperature is 515-535℃; and then obtaining the hot-pressed alloy through deformation hot working; wherein the total deformation reaches 85%-92%.

4. The preparation process of a high-strength alloy material according to claim 1, characterized in that, The preparation of the alloy precursor includes the following steps: solution treatment and quenching of the hot-pressed alloy, followed by cryogenic treatment in an ultra-low temperature cryogenic chamber; obtaining a stretched alloy through plastic pre-strain; performing gradient annealing of the stretched alloy using a high-power induction heating device at an annealing heating rate of 18-25℃ / s, followed by secondary water quenching; and then performing the three-stage aging treatment to obtain the alloy precursor.

5. The preparation process of a high-strength alloy material according to claim 4, characterized in that, The three-stage aging treatment includes the following processes: the first stage is carried out at 110-120℃; the second stage is carried out at 135-145℃; and the third stage is carried out at 160-170℃ to obtain the alloy precursor.

6. The preparation process of a high-strength alloy material according to claim 1, characterized in that, The preparation of the alloy material includes the following steps: cleaning the alloy precursor with anhydrous ethanol, followed by argon plasma cleaning; then introducing a mixed gas of acetylene and hydrogen, with an acetylene flow rate of 15-30 sccm and a hydrogen flow rate of 5-10 sccm, and obtaining a carbon deposition layer with a thickness of 200-500 nm through chemical vapor deposition, followed by furnace cooling under argon protection to obtain the alloy material.

7. A high-strength alloy material, characterized in that, The high-strength alloy material is prepared by the preparation method described in any one of claims 1-6; by mass percentage, it contains 2.2-2.6% lithium, 3.2-3.8% copper, 0.12-0.18% scandium, 0.05-0.09% indium, 0.35-0.45% silver, 0.1-0.15% zirconium, and 0.01-0.05% ytterbium, with the balance being aluminum.

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