Ultrahigh-strength-plasticity advanced high-strength steel based on martensite structure genetic effect and preparation method of ultrahigh-strength-plasticity advanced high-strength steel

Abstract

The invention discloses martensite structure genetic effect-based ultrahigh strength and plasticity advanced high-strength steel and a preparation method thereof, and belongs to the technical field of metal materials. The advanced high-strength steel comprises the following chemical components: C, Mn, Si, Al, Nb, Ti, Fe and the like, and the core of the advanced high-strength steel is that a microstructure is regulated and controlled through the'structure genetic effect 'of martensite: full martensite formed by pre-quenching is used as an initial structure, and reverse transformed austenite is preferentially nucleated on a martensite lath interface through critical zone annealing and inherits the lath form of the martensite lath; the double-peak retained austenite with battens as the main part (the volume fraction is larger than or equal to 70%) and blocks as the auxiliary part (the volume fraction is smaller than or equal to 30%) is finally obtained, the Mn element of the batten retained austenite is obviously enriched, the mass fraction of the Mn element is larger than or equal to 3.5%, the tensile strength of the finally-obtained advanced high-strength steel reaches 840-900 MPa, the ductility is 35-40%, the product of strength and ductility is larger than or equal to 31.5 Pa%, and the tensile strength is larger than or equal to 31.5 Pa%. The method is suitable for the fields of automobile lightweight, high-end equipment manufacturing and the like with strict requirements on material performance.

Description

Technical Field

[0001] This invention relates to an advanced high-strength steel, particularly to an ultra-high strength and ductility advanced high-strength steel based on the "hereditary effect of martensite structure", and also to its preparation method, belonging to the field of metal materials technology. Background Technology

[0002] With the rapid development of new energy vehicles, vehicle lightweighting has become a key approach to reducing energy consumption and increasing driving range. Advanced high-strength steel (AHSS), due to its excellent strength-ductility balance, has become one of the core materials for achieving vehicle weight reduction. In the research and development of AHSS, utilizing the transformation-induced plasticity (TRIP effect) of retained austenite to improve the strength-ductility product of the material is one of the most effective methods currently available. This has led to the development of a series of steel grades such as transformation-induced plasticity steel (TRIP steel), quenched and partitioned steel (Q&P steel), and medium manganese steel.

[0003] However, the performance and processes of existing AHSS still have significant limitations:

[0004] 1. Performance bottleneck: Traditional TRIP steel and Q&P steel are mostly based on ferrite / pearlite as the initial structure. The residual austenite is controlled by critical zone annealing and over-aging treatment. Their strength-ductility product is generally limited to 25~28GPa·%, which is difficult to meet the high plasticity forming requirements of complex parts of new energy vehicles (such as anti-collision beams and battery shells).

[0005] 2. Microstructure defects: The retained austenite obtained by existing processes is mostly blocky, and the carbon content and stability are insufficient. The TRIP effect is triggered too early or not sustained during deformation, making it difficult to improve strength and plasticity in a coordinated manner.

[0006] 3. Technological limitations: To overcome performance bottlenecks, existing technologies often rely on high alloying (such as increasing the Mn content to over 5%) or complex multi-step heat treatment (such as multi-stage quenching and partitioning, multi-step critical zone annealing, etc.), which not only significantly increases production costs but also deteriorates the weldability and formability of materials, limiting industrial applications.

[0007] Therefore, optimizing the morphology, content, and stability of retained austenite through innovative microstructure control mechanisms under conventional composition systems and simplified processes has become the key to breaking through the performance bottleneck of existing AHSS. Existing published literature (“Pre-quenching induced lath structures and enhanced TRIP effect to optimize the strength-ductility of ultrahigh-strength hot-galvanized steel”, Chengyu Guo, et al., journal of materials research and technology, 2023; 26:7958~7971) discloses an advanced high-strength steel that uses pre-quenching to introduce a martensitic matrix for microstructure design. However, it mainly uses a low-silicon, high-alumina composition (Si: 0.85wt.%; Al: 0.91wt.%), which not only easily clogs the crystallizer during production but also significantly increases the pre-quenching temperature (to 950℃). This adversely affects the stability and production cost of industrial production. Furthermore, its over-aging process parameters are specifically designed for hot-dip galvanizing, with a temperature as high as 460℃ and a very short over-aging time of only 30 seconds. In addition, the C content in the retained austenite is low (1.01~1.07wt.%), which leads to poor mechanical stability of the retained austenite. It is prone to phase transformation induced plasticity (TRIP) effect in the early stage of deformation, thus contributing less to plasticity and resulting in lower elongation and strength-ductility product. Summary of the Invention

[0008] To address the technical problems of insufficient strength-ductility product and insufficient mechanical stability of retained austenite in existing advanced high-strength steels, the first objective of this invention is to provide an ultra-high strength-ductility advanced high-strength steel based on the "microstructure inheritance effect" of martensite. The retained austenite in this steel exhibits a bimodal distribution, with lath as the main component (volume fraction ≥70%) and blocky as the secondary component (volume fraction ≤30%). Furthermore, the lath retained austenite shows significant enrichment of Mn element (Mn element mass fraction ≥3.5%). This characteristic endows the metallic material with the advantages of both high strength and high ductility, meeting the demand for ultra-high strength-ductility matching materials in fields such as automotive lightweighting.

[0009] The second objective of this invention is to provide a method for preparing an advanced high-strength steel with ultra-high strength and ductility. This method utilizes the "structure inheritance effect" of the initial martensitic structure to refine the structure, increase the content of residual austenite and carbon, and simultaneously optimize and control its morphology and stability, ultimately obtaining an advanced high-strength steel with ultra-high strength and ductility matching.

[0010] To achieve the above technical objectives, this invention provides an advanced high-strength steel with ultra-high strength and ductility based on the "microstructure inheritance effect" of martensite, comprising the following alloying elements by mass percentage: C 0.18~0.22%; Mn 1.8~2.3%; Si 1.2~1.6%; Al 0.4~0.6%; Nb 0.02~0.04%; Ti 0.05~0.07%; the remainder being Fe and unavoidable impurities; its microstructure comprises critical zone ferrite, carbide-free bainite, and retained austenite; the carbide-free bainite is composed of bainitic ferrite and martensite islands; the retained austenite forms a bimodal distribution with lath as the main component and blocky as the secondary component through the microstructure inheritance effect of martensite; the mass fraction of Mn in the lath-morphological retained austenite is ≥3.5%.

[0011] In the advanced high-strength steel of this invention, carbon (C) is the core element for regulating the stability of retained austenite. It primarily enhances the steel's strength through interstitial solid solution strengthening and significantly improves the thermal and mechanical stability of austenite, delaying its decomposition during cooling and transformation during deformation. Under the "microstructure inheritance effect" of pre-quenching and critical zone annealing, carbon preferentially distributes to the reverse-transformed austenite at the martensitic lath interface, resulting in a carbon content of 1.35~1.5% in the lath retained austenite (XRD verification), significantly improving the mechanical stability of the retained austenite. When the carbon content is too low (<0.18%), the carbon distribution of retained austenite is insufficient, leading to premature transformation into martensite during deformation, resulting in a decrease in elongation and strength-ductility product (e.g., below 25 GPa·%). When the carbon content is too high (>0.22%), coarse martensitic islands are easily formed, increasing brittleness and deteriorating weldability. Simultaneously, the excessive stability of austenite may inhibit the triggering of the TRIP effect. Manganese (Mn) is a strong austenitizing element that can effectively expand the austenite region and synergistically improve the thermal and mechanical stability of the retained austenite with carbon.

[0012] In the advanced high-strength steel of this invention, manganese is significantly enriched in the lath retained austenite (EPMA verification, content ≥3.5%), forming a concentration difference of ≥2.0% at the ferrite phase boundary, further inhibiting the transformation of retained austenite during deformation and prolonging the duration of the TRIP effect. The solid solution strengthening effect of manganese can also improve the matrix strength while delaying the formation of equilibrium phases such as pearlite and bainite, creating a time window for retained austenite. When the manganese content is below 1.8%, the austenite stability is insufficient, and it easily decomposes into bainite during the over-aging stage, with the retained austenite volume fraction dropping below 13%, and the strength-ductility product less than 25 GPa·s. When the manganese content is above 2.3%, compositional segregation easily occurs, leading to temper brittleness and embrittlement of the weld heat-affected zone, while also increasing the degree of cold rolling work hardening and reducing formability.

[0013] The advanced high-strength steel of this invention features high silicon and low aluminum content. The core role of silicon (Si) is to strongly suppress the precipitation of cementite (Fe3C). During over-aging, it prevents carbon from diffusing from ferrite / bainitic ferrite to austenite, thus preventing the formation of carbides and forcing more carbon to remain in austenite, thereby increasing the carbon content and stability of the retained austenite. Silicon's solid solution strengthening effect is relatively weak, with minimal impact on plasticity. In synergy with aluminum, it further enhances the suppression of carbides in carbide-free bainite, ensuring that the martensite island exists in a mixed form of carbon-rich austenite and martensite, avoiding plasticity deterioration caused by brittle carbides. When the silicon content is below 1.2%, the ability to suppress carbides is insufficient, and Fe3C easily precipitates during the over-aging stage, reducing the carbon content of the retained austenite to below 1.2% and its volume fraction to below 8%. When the silicon content is above 1.6%, it leads to the formation of excessively thick iron oxide scale during hot rolling, which is difficult to remove by pickling, affecting the surface quality of the cold-rolled sheet. Aluminum (Al) primarily functions as a deoxidizer, reducing the oxygen content in steel and improving its purity and toughness. Simultaneously, similar to silicon, aluminum helps suppress carbide precipitation, enhances the stability of retained austenite, and optimizes the ratio of martensite to retained austenite by slightly increasing the Ms point (formula verification). In this invention, aluminum and silicon (total content 1.6~2.2%) work together to suppress carbide precipitation during critical annealing and over-aging stages, maintaining the carbon content of retained austenite above 1.35%. When the aluminum content is below 0.4%, deoxidation is insufficient, oxide inclusions increase, and these inclusions easily become crack initiation sites during stretching, leading to a decrease in elongation. When the aluminum content is above 0.6%, the annealing temperature is significantly increased, and the crystallizer is easily blocked during continuous casting.

[0014] In this invention, niobium (Nb) and titanium (Ti) are used as microalloying elements in the advanced high-strength steel. They form nanoscale carbides (NbC, TiC) with carbon to refine grains and enhance yield strength through precipitation strengthening, while significantly inhibiting recrystallization. NbC has low solubility at high temperatures (≥1000℃), allowing it to precipitate early during hot rolling and pin austenite grain boundaries, thus inhibiting grain growth during hot rolling. TiC is more stable at medium temperatures (600~800℃) and precipitates during the critical annealing stage, preventing the coarsening of austenite grains due to reverse transformation. When the Nb content is <0.02% or the Ti content is <0.05%, the carbide precipitation is insufficient, resulting in poor grain refinement and a yield strength <600MPa. When Nb >0.04% or Ti >0.07%, the alloy cost increases significantly, and the grain refinement and strengthening effects reach saturation and cannot be further improved.

[0015] The advanced high-strength steel of this invention comprises a microstructure consisting of critical zone ferrite, carbide-free bainite, and retained austenite. Critical zone ferrite, the "soft phase," possesses relatively low strength and good plasticity, primarily enhancing the material's plasticity and formability. As one of the matrix phases, critical zone ferrite alleviates stress concentration through its deformation coordination ability, providing deformation space for the subsequent TRIP effect of retained austenite. Carbide-free bainite, composed of bainitic ferrite and Mao islands, mainly contributes to the balance between strength and plasticity. Bainitic ferrite exhibits certain strength and good toughness; its lath-like morphology guides stress transmission, preventing stress concentration at single-phase boundaries. Mao islands (a mixture of retained austenite and martensite), as a hard phase, exert a strengthening effect through dispersed distribution, enhancing the overall strength of the material. The key advantage of carbide-free bainite lies in avoiding the brittle precipitation of carbides in traditional bainite. Its interface with ferrite and retained austenite can serve as a trigger site for the TRIP effect, coordinating the deformation of each phase during deformation and delaying necking. Retained austenite is the "core functional phase" for achieving ultra-high strength-ductility matching, significantly improving the material's plasticity and work hardening capacity through the TRIP effect. In the advanced high-strength steel of this invention, retained austenite exists in a bimodal distribution, primarily lath-like with some blocky components: lath-like retained austenite is oriented along the original martensite lath interface, exhibiting high carbon content and high Mn enrichment (≥3.5%), excellent stability, and can continue to transform into martensite under large strains, providing sustained work hardening; blocky retained austenite undergoes phase transformation first in the early stages of deformation, rapidly increasing the work hardening rate and avoiding early plastic instability.

[0016] As a preferred embodiment, the retained austenite accounts for 15-20% of the total volume of the alloy, and the carbon mass content is 1.35-1.5%.

[0017] As a preferred embodiment, the volume percentage of lath-shaped retained austenite in the retained austenite is ≥70%, and the volume percentage of blocky retained austenite is ≤30%. The synergistic effect of the appropriately proportioned lath-shaped and blocky retained austenite enables the material to achieve a tensile strength of 840~900MPa while maintaining an elongation of 35~40%, ultimately achieving a strength-ductility product of over 31.5GPa·%.

[0018] The advanced high-strength steel of this invention features a bimodal distribution of retained austenite, with lath as the main component and block as the secondary component (laminate accounting for ≥70%). The lath retained austenite is oriented along the original martensite lath interface, and Mn element is enriched at the phase boundary (concentration difference ≥2.0%), which significantly improves the mechanical stability of retained austenite and the persistence of the TRIP effect.

[0019] In this invention, "strip" usually refers to a length axis ratio of 2 or more, and "block" usually refers to a length axis ratio of less than 2, in order to distinguish the morphology of residual austenite.

[0020] As a preferred embodiment, the average size of the Mao Island is 2~3μm or less than 1μm.

[0021] As a preferred embodiment, the volume of the critical region ferrite accounts for 30-40% of the total alloy volume. In the advanced high-strength steel of the present invention, the volume fraction of critical region ferrite is stably maintained at 30-40%. Its uniformly distributed fine grains (average size 2-3 μm) can avoid early fracture caused by excessive local deformation. At the same time, through interfacial interaction with hard phases (such as Mao Island, with an average size of 2-3 μm or less than 1 μm), the overall microstructure is refined, indirectly improving the work hardening ability of the material.

[0022] As a preferred embodiment, the microstructure also includes primary martensite, primarily because primary martensite is generated when the temperature of primary overaging is below the Ms point.

[0023] This invention also provides a method for preparing ultra-high strength and ductility advanced high-strength steel based on the "hereditary effect of martensite structure". The method involves melting, casting and forging alloying elements to form a steel billet; the steel billet is then subjected to homogenization treatment, hot rolling treatment, cold rolling treatment, pre-quenching treatment, critical zone annealing treatment and one-step aging treatment to obtain the final product.

[0024] As a preferred embodiment, the homogenization treatment is performed at 1230~1260℃ for 3~5 hours. Homogenization promotes alloy element diffusion through high-temperature holding, eliminating dendritic segregation and compositional inhomogeneity in the ingot, laying the foundation for the uniformity of the microstructure in subsequent processes. The temperature and holding time significantly affect the effect. 1230~1260℃ is the optimal range balancing diffusion efficiency and grain control: at this temperature, the billet is fully austenitized, atomic diffusion is active, and microscopic segregation of elements such as C and Mn can be effectively eliminated. If the temperature is too low (<1230℃), incomplete austenitization will occur, insufficient diffusion of elements such as Nb and Ti, difficulty in eliminating dendritic segregation, and subsequent formation of banded structures, leading to anisotropy in mechanical properties. If the temperature is too high (>1260℃), abnormal coarsening of austenite grains will occur, increasing the difficulty of hot rolling refinement, and even leading to grain boundary weakening and cracking, while also exacerbating surface oxidation and burn-off. The optimal holding time is 3 to 5 hours to ensure sufficient diffusion of elements. When the holding time reaches 3 hours, significant segregation can be basically eliminated. When the holding time reaches 5 hours or more, the uniformity can be further improved by 10 to 15%. However, when the holding time exceeds 5 hours, the diffusion driving force becomes saturated, and the improvement in effect is limited. On the contrary, the long holding time leads to coarse grains and increases costs.

[0025] As a preferred embodiment, the hot rolling process is performed under the following conditions: hot rolling to a thickness of 3.8~4.2 mm.

[0026] As a preferred embodiment, the cold rolling process is performed under the following conditions: cold rolling to a thickness of 1.2~1.6 mm.

[0027] As a preferred embodiment, the pre-quenching treatment conditions are: annealing at 860~900℃ for 300~500s, followed by water quenching to room temperature. Pre-quenching yields a fully martensitic microstructure with a lath bundle width of 0.5~1μm, serving as a "microstructure genetic carrier." Pre-quenching is the core step in constructing the martensitic "microstructure genetic carrier." The process parameters of annealing at 860~900℃ for 300~500s followed by water quenching to room temperature allow for precise control of the refinement and integrity of the fully martensitic microstructure. This temperature range ensures complete austenitization of the billet and uniform austenite grains (size controlled at 20~30μm), providing a uniform matrix for subsequent martensitic transformation. The 300~500s holding time balances the uniformity of austenite composition with the risk of grain coarsening, avoiding compositional segregation due to insufficient holding or grain coarsening due to excessive holding. Water quenching can suppress the precipitation of equilibrium phases such as pearlite and bainite, ensuring the formation of a fully martensitic microstructure. This not only provides a sufficient interface for the nucleation of austenite in the subsequent critical zone annealing but also guides the newly formed austenite to inherit the lath morphology through the "microstructure inheritance effect," laying the foundation for the final formation of a bimodal residual austenite with "lath as the main component and blocky as the secondary component." If the temperature is below 860℃, austenitization is incomplete, leaving undissolved carbides or ferrite, which disrupts the continuity of the martensitic microstructure and weakens the inheritance effect. If the temperature is above 900℃, the austenite grains coarsen. If the holding time is too short (<300s), the uneven martensitic transformation will result due to the inhomogeneity of the austenite composition. If the holding time is too long (>500s), it will exacerbate grain coarsening, both of which are detrimental to subsequent microstructure control.

[0028] As a preferred embodiment, the critical zone annealing treatment is performed at 800-830℃ for 280-350 seconds. The martensitic lath interface serves as the preferential nucleation site for reverse austenite transformation, ensuring the newly formed austenite inherits the lath morphology.

[0029] As a preferred solution, the conditions for the one-step over-aging treatment are as follows: at a temperature of (Ms - 30°C) to (Ms + 60°C), after holding for 500 to 650 s, air cooling; where Ms is the starting temperature of martensitic transformation Ms, and Ms = 539 - 423×w(C) - 30.4×w(Mn) - 7.5×w(Si) + 30×w(Al), and w(C), w(Mn), w(Si), and w(Al) are the mass fractions of the corresponding elements respectively. After the one-step over-aging treatment, the lath retained austenite is distributed directionally along the original martensite lath interface, and the enrichment concentration difference of Mn element at the phase boundary between the lath retained austenite and ferrite is ≥2.0% (mass fraction). The one-step over-aging treatment is the key link to regulate the morphology, stability, and element distribution of the retained austenite. The temperature range of (Ms - 30°C) to (Ms + 60°C) and the holding time of 500 to 650 s directly determine the final manifestation of the "tissue genetic effect". This temperature range precisely matches the thermodynamic equilibrium between martensite and austenite: when the temperature is lower than the Ms point, part of the formed martensite can increase the carbon content of the retained austenite through carbon atom partitioning; when the temperature is higher than the Ms point, the stability of austenite is sufficient to resist martensitic transformation, and at the same time, it promotes the enrichment of Mn element to the lath retained austenite / ferrite phase boundary (concentration difference ≥2.0%). The holding of 500 to 650 s provides sufficient time for element diffusion, and the carbon content is stabilized at 1.35 to 1.5%. If the temperature is too low (<Ms - 30°C), a large amount of austenite will be transformed into primary martensite, resulting in a reduction in the amount of lath retained austenite (<8%) and insufficient Mn enrichment (<1.0%), and the TRIP effect is weakened; if the temperature is too high (>Ms + 60°C), the austenite is easily decomposed into bainite or carbide precipitates, the carbon content drops below 1.0%, the lath proportion decreases (<50%), and the plasticity deteriorates (elongation <20%). If the holding time is too short (<500 s), the element diffusion is insufficient and the stability of austenite is insufficient; if the holding time is too long (>650 s), a large amount of decomposition of austenite and softening of the matrix will occur, which is not conducive to the improvement of the strength-ductility product. This process balances the thermodynamics of phase transformation and the kinetics of diffusion to ensure the optimal morphology, content, and element distribution of the lath retained austenite, and realizes the synergistic effect of strength and plasticity.

[0030] A preparation method of an ultra-high strength and plasticity advanced high-strength steel based on the martensite "tissue genetic effect" of the present invention includes the following steps:

[0031] (1) Alloy raw materials are proportioned according to the chemical composition and stoichiometry of alloy elements, and the alloy element raw materials are melted into an ingot, and then forged into a steel billet;

[0032] (2) The steel billet is heated to 1230 to 1260°C, held for 3 to 5 h for homogenization treatment, and then hot-rolled to 3.8 to 4.2 mm, and immediately wrapped with asbestos after hot-rolling to simulate the coiling process;

[0033] (3) Hot-rolled steel plates are pickled with HCl with a volume fraction of 18-22% and then cold-rolled to 1.2-1.6 mm;

[0034] (4) Pre-quenching treatment: The cold-rolled steel bar is fully austenitized and annealed at 860~900℃ for 300~500s and then water-quenched to room temperature to obtain a full martensitic structure with a strip bundle width of 0.5~1μm;

[0035] (5) Critical zone annealing: The pre-quenched steel plate is held at 800~830℃ for 280~350s, and the martensitic lath interface is used as the preferential nucleation site for reverse transformation of austenite to inherit its lath morphology.

[0036] (6) One-step over-aging treatment: After the steel plate annealed in the critical zone is held at (Ms-30℃)~(Ms+60℃) for 500~650s and then air-cooled; wherein, the martensitic transformation start temperature Ms is calculated according to the formula "Ms=539-423×w(C)-30.4×w(Mn)-7.5×w(Si)+30×w(Al)" (where w(C), w(Mn), w(Si), and w(Al) are the mass fractions of the corresponding elements respectively). After the one-step over-aging treatment, the lath residual austenite is oriented along the original martensitic lath interface, and the enrichment concentration difference of Mn element at the phase boundary between the lath residual austenite and ferrite is ≥2.0% (mass fraction).

[0037] The ultra-high strength and ductility advanced high-strength steel of this invention, based on the designed alloy element composition of C (0.18~0.22%), Mn (1.8~2.3%), Si (1.2~1.6%), Al (0.4~0.6%), Nb (0.02~0.04%), and Ti (0.05~0.07%), is processed through melting, forging, hot rolling, cold rolling, pre-quenching (full austenitization followed by annealing and water quenching at 860~900℃), critical zone annealing (800~830℃), and one-step over-aging (Ms-30℃~Ms+60℃) to control the formation of the microstructure of the advanced high-strength steel. Through the martensitic "structure inheritance effect," it regulates the formation of an alloy phase mainly composed of critical zone ferrite, carbide-free bainite (containing bainitic ferrite and martensitic islands), and retained austenite (at the over-aging temperature). Below the Ms point, it contains primary martensite, and the retained austenite exhibits a bimodal distribution of "mainly lath and supplemented by blocky". The volume fraction of lath morphology (length-to-width ratio ≥3) is ≥70%, and the volume fraction of blocky morphology is ≤30%. The mass fraction of Mn element in the lath retained austenite is ≥3.5%, the total integral of retained austenite is 15~20%, and the carbon content is 1.35~1.5%. The resulting advanced high-strength steel has a tensile strength of 840~900MPa, an elongation of 35~40%, and a strength-ductility product ≥31.5GPa·.

[0038] Compared with the prior art, the technical solution of the present invention has the following beneficial effects:

[0039] 1. Organizational Innovation: Utilizing the "organic genetic effect" of martensite, a bimodal distribution of retained austenite is achieved, characterized by "lath as the main component and block as the secondary component" (lath accounting for ≥70%). The lath retained austenite is oriented along the original martensite lath interface, and Mn element is enriched at the phase boundary (concentration difference ≥2.0%), which significantly improves the mechanical stability of retained austenite and the persistence of the TRIP effect.

[0040] 2. Performance breakthrough: The resulting steel has a tensile strength of 840~900MPa, an elongation of 35~40%, and a strength-ductility product ≥31.5GPa·%, with the optimal embodiment reaching a strength-ductility product of 32.6GPa·%, far exceeding the performance levels of traditional TRIP steel and Q&P steel;

[0041] 3. Process controllability: Through the synergistic process of pre-quenching-critical zone annealing-one-step over-aging, the "microstructure inheritance effect" can be precisely controlled. The process parameter range is clear, the repeatability is strong, and it is suitable for industrial production.

[0042] 4. Wide range of applications: The ultra-high strength and plasticity matching characteristics of advanced high-strength steel materials make them suitable for complex molded parts such as anti-collision beams and chassis components of new energy vehicles, meeting the dual requirements of lightweighting and safety. Attached Figure Description

[0043] Figure 1 This is a flowchart of the heat treatment process of the present invention.

[0044] Figure 2 This is a SEM image of a typical microstructure in the advanced high-strength steel of Embodiment 1 of the present invention.

[0045] Figure 3 The results are typical EPMA elemental analysis results in the advanced high-strength steel of Example 1 of this invention.

[0046] Figure 4 The tensile curves are those of the advanced high-strength steel in Examples 1-2 of this invention. Detailed Implementation

[0047] To further illustrate the content of this invention, the following detailed description is provided in conjunction with specific embodiments. The scope of protection of the claims is not limited by the following specific embodiments.

[0048] In the following examples and comparative examples, all percentages are weight percentages.

[0049] Examples 1-2

[0050] Based on the composition in Table 1, the ingot was smelted and cast into a billet, which was then forged to obtain a 60×60mm cross-section billet. The billet was then hot-rolled according to the following process: the billet was heated to 1250℃ and held for 4 hours, followed by 7 passes of rolling to obtain a 4.5mm thick hot-rolled plate. The initial rolling temperature was approximately 1140℃, and the final rolling temperature was approximately 890℃. After hot rolling, asbestos was wrapped to simulate coiling, and the final hot-rolled plate had a microstructure of ferrite + pearlite. The hot-rolled plate was then pickled with 20% hydrochloric acid and cold-rolled to obtain a 1.4mm thick cold-rolled plate.

[0051] The heat treatment process of the embodiment is as follows Figure 1 The specific parameters are shown in Table 2: First, the cold-rolled sheet is heated to 860~900℃, held for 300~500s, and then water-quenched to room temperature to obtain a fully martensitic structure; then, the steel sheet is reheated to 800~820℃, held for 300~350s, and then salt-bath quenched to 375~400℃, held for 560~600s, and subsequently air-cooled to room temperature. The final microstructure and elemental properties under different heat treatment process parameters are as follows: Figures 2-3 As shown, the corresponding tissue characteristics and mechanical properties are shown in Table 3 and Figure 4 As shown.

[0052] Comparative Examples 1-4

[0053] The composition of the comparative example is shown in Table 1, the heat treatment process parameters are shown in Table 2, the heating and cooling methods used are the same as in Example 1, and the corresponding microstructure and mechanical properties are shown in Table 3.

[0054]

[0055]

[0056]

[0057] According to Tables 2 and 3:

[0058] Comparative Example 1, lacking pre-quenching treatment and martensitic lath structure as a "microstructure genetic carrier," suffers from the inefficiency of carbon and manganese distribution along the lath interface, resulting in a significant reduction in the content of retained austenite and carbon, a substantial decrease in the proportion of retained austenite in the lath, insufficient contribution of the TRIP effect, and ultimately low elongation and strength-ductility product.

[0059] The overaging temperature of Comparative Example 2 was too low. A large amount of austenite was transformed into martensite at the overaging temperature, and the remaining retainable austenite content was reduced. The ability of the TRIP effect to persist was weakened, resulting in lower elongation and strength-ductility product.

[0060] The pre-quenching temperature of Comparative Example 3 was too low, placing it in the two-phase region and resulting in a mixed structure of equiaxed ferrite and martensite. This disrupted the continuity and content of the lath structure, reduced the subsequent austenite nucleation sites, decreased the content of retained austenite, and resulted in insufficient lath proportion. Consequently, the TRIP effect contributed only to a limited extent, and both the elongation and strength-ductility product were low.

[0061] The critical annealing temperature of Comparative Example 4 was too high, and it had entered the full austenitic region. The martensitic lath structure formed by pre-quenching completely disappeared, the "hereditary effect of microstructure" was eliminated, the retained austenite was mainly blocky and its content was significantly reduced. The TRIP effect contributed insufficiently, and the final elongation and strength-ductility product were both low.

Claims

1. An advanced high-strength steel with ultra-high strength and ductility based on the "martensitic structure inheritance effect", characterized in that: It contains the following alloying elements by mass percentage: C 0.18~0.22%; Mn 1.8~2.3%; Si 1.2~1.6%; Al 0.4~0.6%; Nb 0.02~0.04%; Ti 0.05~0.07%; the remainder being Fe and unavoidable impurities; Its microstructure consists of critical region ferrite, carbide-free bainite, and retained austenite; The carbide-free bainite is composed of bainitic ferrite and Mao islands; The retained austenite forms a bimodal distribution with laths as the main component and blocky components as the secondary component through the tissue genetic effect of martensite. The mass fraction of Mn in the lath-shaped residual austenite is ≥3.5%.

2. The advanced high-strength steel based on the "martensitic structure inheritance effect" according to claim 1, characterized in that: The retained austenite accounts for 15-20% of the total volume of the alloy, and its carbon mass content is 1.35-1.5%.

3. An advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 1 or 2, characterized in that: The volume percentage of lath-shaped retained austenite is ≥70%, and the volume percentage of blocky retained austenite is ≤30%.

4. The advanced high-strength steel based on the "martensitic structure inheritance effect" according to claim 1, characterized in that: The average size of the Mao Island is 2~3μm or less than 1μm; The volume of ferrite in the critical region accounts for 30-40% of the total volume of the alloy.

5. A method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" as described in any one of claims 1 to 4, characterized in that: Alloying raw materials are smelted, cast into ingots and forged to form steel billets; the steel billets are then subjected to homogenization treatment, hot rolling treatment, cold rolling treatment, pre-quenching treatment, critical zone annealing treatment and one-step aging treatment to obtain the final product.

6. The method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 5, characterized in that: The homogenization treatment is performed at a temperature of 1230~1260℃ for 3~5 hours.

7. The method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 5, characterized in that: The hot rolling conditions are as follows: hot rolling to a thickness of 3.8~4.2 mm; The conditions for the cold rolling process are: cold rolling to a thickness of 1.2~1.6mm.

8. The method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 5, characterized in that: The pre-quenching treatment conditions are as follows: annealing at 860~900℃ for 300~500s, followed by water quenching to room temperature.

9. The method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 5, characterized in that: The critical zone annealing conditions are: holding at 800~830℃ for 280~350s.

10. The method for preparing an ultra-high strength and ductile advanced high-strength steel based on the "martensitic structure genetic effect" according to claim 5, characterized in that: The conditions for the one-step over-aging treatment are as follows: after holding at a temperature of (Ms-30℃) to (Ms+60℃) for 500-650s, air-cooled; where Ms is the martensitic phase transformation start temperature Ms, Ms=539-423×w(C)-30.4×w(Mn)-7.5×w(Si)+30×w(Al), and w(C), w(Mn), w(Si), and w(Al) are the mass fractions of the corresponding elements.

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