2300mpa grade high strength and toughness hot forming steel and method of making the same

Abstract

The application discloses a 2300MPa-grade high-strength and high-toughness hot forming steel and a preparation method thereof, and comprises the following chemical components in percentage by mass: C: 0.46-0.52%, Si: 0.30-0.50%, Mn: 0.40-1.00%, Cr: 0.15-0.50%, B: 0.001-0.004%, Al: 0.15-0.50%, V: 0.10-0.20%, Nb: 0.02-0.05%, and N≤0.005%; and 4≤V / Nb≤8, 80≤Al / N≤120, and V, Nb, Al and N represent the mass percentage content of corresponding elements. The content of each component is determined based on strengthening contribution distribution and hierarchical precipitation control, and the obtained steel product can simultaneously meet the requirements of more than 2300MPa ultra-high strength, good plasticity, excellent cold bending property, low hydrogen embrittlement sensitivity and cost effectiveness.

Description

Technical Field

[0001] This invention relates to the field of ultra-high strength automotive steel technology, and in particular to a 2300MPa grade high-strength and high-toughness hot-formed steel and its preparation method. Background Technology

[0002] Hot-formed steel (PHS) has become the preferred material for manufacturing critical collision safety components such as A-pillars, B-pillars, and door anti-collision beams due to its excellent properties such as ultra-high strength, low springback, and high forming precision. In recent years, its strength grades have undergone rapid iteration, moving from the mainstream 1500MPa level to 1800MPa, 2000MPa, and even higher strength levels. However, the leap in strength has not simultaneously resolved the inherent contradiction between strength and toughness in the material; instead, it has brought more severe challenges: on the one hand, while strength increases, plasticity and bending properties often deteriorate sharply; on the other hand, with the leap in strength, the material's sensitivity to hydrogen-induced delayed fracture (hydrogen embrittlement) increases significantly, which seriously threatens the safety of vehicles in long-term service or in corrosive environments, becoming one of the core bottlenecks limiting its application.

[0003] To address these challenges, existing technologies have mainly formed three mainstream routes, but all have significant inherent drawbacks: First, the high-carbon, high-alloy route. This relies on increasing carbon content and adding large amounts of precious elements such as Mo and Ni. For example, Chinese invention application CN202311430074.4 discloses a medium-thickness 2200MPa grade ultra-high strength steel and its preparation method, using a FeCSiMnCrNiMo composition system. This results in high material costs, and the interaction between interstitial solid-solid carbon atoms and dislocations forms strong traps, competing with hydrogen atoms. This may actually promote the enrichment of hydrogen in weak regions such as grain boundaries, worsening the resistance to hydrogen embrittlement. Second, the transformation-induced plastic toughening route based on retained austenite. For example, Chinese invention patent application CN202210703407.5 discloses a high-plasticity hot-formed steel and its rapid heating preparation method and application. It employs a multiphase microstructure composed of a martensitic matrix, submicron retained austenite, and nano-carbides. While the retained austenite plasticizing route can improve ductility, its strength upper limit is usually insufficient. Furthermore, the content, distribution, and stability of retained austenite are extremely sensitive to composition and heat treatment parameters, resulting in poor stability in large-scale production. In addition, metastable austenite can either act as a reversible hydrogen trap to capture hydrogen or undergo phase transformation under specific conditions to become a hydrogen diffusion channel or release source, complicating the assessment and control of hydrogen embrittlement risk. Thirdly, there is the Ti-based microalloying composite strengthening route. For example, Chinese invention patent application CN201710030911.2 discloses a method for preparing hot-formed steel with a tensile strength ≥2000MPa. This route adopts a "medium carbon + Ti, Nb composite microalloying" design. Its core logic is to add Ti to fix nitrogen in the molten steel, ensuring the effective solid solubility of B. However, during the solidification or high-temperature homogenization of molten steel, coarse, angular TiN inclusions with a size of micrometers (or even exceeding 10 μm) are easily formed. These non-deformable, fragile, hard, and brittle phases bonded to the matrix not only act as crack initiation sites, impairing bending and fatigue properties, but their interface with the matrix may also become a hydrogen accumulation zone, worsening resistance to hydrogen embrittlement.

[0004] In summary, existing technologies exhibit irreconcilable contradictions and significant shortcomings across multiple performance dimensions, including cost, strength, plasticity, and resistance to hydrogen embrittlement. Therefore, developing a new generation of hot-formed steel that simultaneously meets the requirements of ultra-high strength (above 2300 MPa), good plasticity, excellent cold bending properties, low hydrogen embrittlement sensitivity, and cost-effectiveness has become a key technical challenge for overcoming industry bottlenecks and ensuring the passive safety and service reliability of automobiles. Summary of the Invention

[0005] To overcome the shortcomings of the above-mentioned technologies, the purpose of this invention is to provide a 2300MPa grade high-strength and high-toughness hot-formed steel and its preparation method, which solves the problem that it is difficult to reconcile multiple performance dimensions such as cost, strength, plasticity, and hydrogen embrittlement resistance when using existing technical routes for 2300MPa grade hot-formed steel. By using theoretical models to guide carbon content, optimizing aluminum nitrogen fixation toughening, and coordinating vanadium / niobium (V / Nb) graded precipitation and hydrogen trapping engineering, the invention achieves high toughness and excellent hydrogen embrittlement resistance in hot-formed steel products with strength greater than 2300MPa.

[0006] To achieve the above objectives, the technical solution adopted by the present invention is as follows: A 2300MPa grade high-strength and high-toughness hot-formed steel is characterized by the following chemical composition by mass percentage: C: 0.46-0.52%, Si: 0.30-0.50%, Mn: 0.40-1.00%, Cr: 0.15-0.50%, B: 0.001-0.004%, Al: 0.15-0.50%, V: 0.10-0.20%, Nb: 0.020-0.050%, N≤0.005%; wherein: 4≤V / Nb≤8, 80≤Al / N≤120, and V, Nb, Al, and N represent the mass percentage of the corresponding elements.

[0007] As a preferred embodiment, the C content in the high-strength and high-toughness hot-formed steel satisfies the following formula: C design =[ target -( Fe + + pre + gb )] / (k c ×ρ M ); In the formula, C design The carbon content is designed in wt.%; target The target tensile strength is expressed in MPa. Fe denoted as the lattice strength of pure iron, in MPa; k1, k2, k3, and k4 are the strengthening coefficients of the substitutional solid solution elements Si, Mn, Cr, and Ni, respectively, in MPa / wt.%; X1, X2, X3, and X4 are the mass percentages of the substitutional solid solution elements Si, Mn, Cr, and Ni, respectively, in wt.%; where Ni is an optional element, and X4 is 0 when Ni is absent. pre The contribution to precipitation enhancement is expressed in MPa and is provided by nano-VC. gbThe contribution to grain refinement, measured in MPa, is mainly provided by AlN and Nb (C,N) grain refinement; k c ρ is the interstitial solid solution strengthening coefficient of carbon, expressed in MPa / wt.%; M This is the retention coefficient of carbon in the solid solution after the martensitic phase transformation, and it is dimensionless. This formula is a carbon content calculation model constructed by integrating several classical theories in materials science. Its form does not exist in isolation, but is a simplified engineering model based on specific composition ranges and process conditions.

[0008] As a preferred embodiment, the Fe The value range is 180-220 MPa; the value ranges of k1, k2, k3, and k4 are 45-55, 28-32, 18-22, and 8-12 MPa / wt.%, respectively. pre =α MPa, where α is the coefficient of precipitation strengthening contribution, and the value of α ranges from 330 to 370 MPa / , This represents the mass percentage of V, in wt.%; gb The value range is 80-120 MPa; k c The value range is 4800-5200 MPa / wt.%; ρ M The value range is 0.75-0.85.

[0009] Furthermore, the aforementioned Fe The value is set to 200 MPa; k1, k2, k3, and k4 are set to 50, 30, 20, and 10 MPa / wt.%, respectively; α is set to 350 MPa / , pre =350 MPa This represents the mass percentage of V, in wt.%; gb The value is 100 MPa; k c The pressure is 5000 MPa / wt.%; because this invention designs a method to capture some carbon atoms through VC precipitation, ρ M The preferred value is 0.8, which aims to reduce the effective solid carbon concentration.

[0010] As a preferred embodiment, the chemical composition of the high-strength and high-toughness hot-formed steel further includes: Ti: 0.01-0.03%, Ni≤0.5%, P≤0.015%, S≤0.010%, with the balance being Fe and unavoidable impurities.

[0011] As a preferred embodiment, the 2300MPa grade high-strength and high-toughness hot-formed steel has a tensile strength ≥2300MPa, an elongation after fracture ≥5%, and an ultimate cold bending angle ≥45°.

[0012] The method for preparing the above-mentioned 2300MPa grade high-strength and high-toughness hot-formed steel is characterized by the following steps: Smelting; the smelting process is carried out according to the target composition and its designed content, and the amount of Al added is determined by measuring the content of N element to ensure that the final composition meets 80≤Al / N≤120, and then cast into ingots; in the target composition, the designed content of C meets: C design =[ target -( Fe + + pre + gb )] / (k c ×ρ M ); Hot-rolled; Cold rolling and annealing; Hot forming and quenching are used to obtain the high-strength and high-toughness hot-formed steel.

[0013] As a preferred embodiment, during the hot rolling process, the ingot is first heated to 1150-1250℃ at a heating rate of 5-10℃ / min and held at that temperature for 2-4 hours; then, multiple hot rolling passes are performed, with a total hot rolling reduction rate ≥90% and a final rolling temperature ≥950℃; after hot rolling, the ingot is rapidly cooled to 550-650℃ at a cooling rate greater than or equal to 15℃ / s and held at that temperature for 1-3 hours, finally cooled in the furnace to obtain a hot-rolled plate with a thickness of 2-5mm. The ingot thickness is 50-300mm.

[0014] Furthermore, the cooling rate after hot rolling is preferably ≥20℃ / s, and even more preferably 20-50℃ / s.

[0015] As a preferred embodiment, the reduction rate during the cold rolling process is 30-60%.

[0016] As a preferred embodiment, during the annealing process, the cold-rolled sheet obtained by cold rolling is first heated to 600-800℃ at a heating rate of 5-10℃ / s, held at that temperature for 3-10 minutes, and then cooled in the furnace to obtain the steel sheet.

[0017] As a preferred embodiment, in the hot forming process, the annealed steel sheet is heated to 900-950°C at a heating rate of 10°C / s or higher, held at that temperature for 3-5 minutes, and then hot-formed by in-mold stamping. The preferred heating rate is 10-100°C / s.

[0018] As a preferred embodiment, the quenching process includes directly quenching and cooling the hot-formed steel plate to obtain the high-strength and high-toughness hot-formed steel; the cooling rate of the quenching and cooling process is ≥20℃ / s.

[0019] As a preferred embodiment, the method for preparing the above-mentioned high-strength and high-toughness hot-formed steel includes the following steps: Smelting: Smelting is carried out according to the target composition and its designed content. The amount of Al added is determined by measuring the content of N element, ensuring that the final composition meets the requirement of 80≤Al / N≤120; casting is then performed into ingots; the designed content of C in the target composition meets the following requirement: C design =[ target -( Fe + + pre + gb )] / (k c ×ρ M ); Hot rolling: The ingot is heated to 1150-1250℃ at a heating rate of 5-10℃ / min and held for 2-4 hours to allow the alloying elements (especially Nb) to be fully dissolved; then, multiple hot rolling passes are performed, with a total hot rolling reduction rate ≥90% and a final rolling temperature ≥950℃. This stage promotes the precipitation of AlN and Nb(C,N), pins grain boundaries, and inhibits excessively coarse grains; after hot rolling, the ingot is rapidly cooled to 550-650℃ at a cooling rate of ≥15℃ / s and held for 1-3 hours, and finally cooled in the furnace to obtain a hot-rolled plate with a thickness of 2-5mm; the cooling rate after hot rolling is preferably ≥20℃ / s, and more preferably 20-50℃ / s; Cold rolling and annealing: After pickling, hot-rolled plates are cold-rolled to obtain cold-rolled plates with a reduction rate of 30-60%. After cold rolling, the cold-rolled plates are heated to 600-800℃ at a heating rate of 5-10℃ / s, held for 3-10 minutes, and then cooled in the furnace to obtain steel plates. This process aims to eliminate residual stress during cold rolling and promote the precipitation of VC.

[0020] Hot forming and quenching: The annealed steel sheet is heated to 900-950℃ at a heating rate of not less than 10℃ / s and held for 3-5 minutes; after in-mold stamping, it is rapidly quenched with a quenching cooling rate ≥20℃ / s, aiming to obtain a lath martensitic matrix and ultra-high strength hot-formed steel. The preferred heating rate is 10-100℃ / s.

[0021] The main elements and their uses in this invention are described below: C: To ensure targetThe design also allows for process fluctuations while controlling the effective dissolved carbon concentration to maintain toughness, with the C content designed to range from 0.46% to 0.52%. This design achieves partial carbon diversion through VC precipitation, reducing the embrittlement effect of dissolved carbon while ensuring the overall strength contribution.

[0022] Vanadium (V) and Niobium (Nb): Sufficient V ensures the dispersion and precipitation of high-density nano-VC within the martensitic laths after quenching. Its functions include: ① precipitation strengthening: providing an increase in core strength; ② hydrogen trapping effect: forming a large number of dispersed irreversible hydrogen trap networks. Appropriate amounts of Nb form stable, undissolved Nb(C,N) particles at high temperatures. Its functions include: ① grain refinement strengthening: effectively pinning austenite grain boundaries and refining grains; ② hydrogen trapping effect: acting as an effective irreversible hydrogen trap to capture hydrogen atoms. This invention achieves "spatiotemporal separation" and synergistic enhancement of the two carbonitrides in terms of precipitation sequence, strengthening function, and hydrogen trapping properties by precisely controlling the V / Nb mass ratio between 4 and 8. When the V content is significantly higher than the Nb content, although it is beneficial for VC precipitation strengthening, it leads to two problems: First, excess V may remain in the solid solution due to incomplete precipitation, adversely affecting hardenability and potentially slightly lowering the Ms point. Second, and more critically, the relatively insufficient Nb content significantly reduces the number of undissolved Nb(C,N) particles formed at high temperatures, weakening the grain refinement effect and resulting in an insufficient number of an important type of irreversible hydrogen trap, which is detrimental to optimizing hydrogen embrittlement resistance. Under this ratio, the material's strength-toughness balance and hydrogen embrittlement resistance are not optimal. When the Nb content is relatively high, although the grain refinement effect is enhanced, it brings new problems: First, the relatively insufficient V content reduces the number of nano-VC particles available for low-temperature precipitation, decreasing the contribution of precipitation strengthening and potentially preventing the strength from reaching the 2300 MPa level target. Secondly, excessive Nb may form relatively large Nb(C,N) particles at high temperatures. If these particles grow excessively, their grain boundary pinning effect may saturate or even reverse, promoting abnormal grain growth. Furthermore, the large particles themselves may become microcrack initiation points, impairing toughness. Simultaneously, economic efficiency deteriorates due to excessive Nb usage. Precisely controlling the V / Nb mass ratio between 4 and 8 achieves a perfect balance between function and cost. This ratio ensures a balanced contribution from precipitation strengthening (VC-dominated) and grain refinement strengthening (NbC-assisted), and constructs a high-density, multi-scale hydrogen trapping system, simultaneously achieving ultra-high strength, good toughness, and excellent resistance to hydrogen embrittlement.

[0023] Al: This invention abandons the traditional Ti nitrogen fixation method. By controlling the Al content to 0.15-0.50% and strictly optimizing the Al / N ratio between 80 and 120, it aims to achieve the best purification and toughening effect. Since N is practically impossible to eliminate in actual production, and even difficult to stably control below 0.002%, this invention uses Al for nitrogen fixation. If the Al content is insufficient, it is impossible to effectively fix all nitrogen atoms in the steel. The residual nitrogen may combine with trace amounts of Ti or other elements in the steel, still posing a risk of forming coarse and harmful nitrides. These inclusions not only impair mechanical properties, but their interfaces are also prone to becoming hydrogen enrichment zones, increasing the risk of hydrogen embrittlement. Excessive Al, while able to fully fix nitrogen, leads to an increase in the number of Al2O3 inclusions in the molten steel, deteriorating the castability of the steel and the cleanliness of the product. At the same time, excessively high Al content will significantly increase the Ac3 temperature, increasing the difficulty and cost of heat treatment process control. By strictly optimizing the Al / N ratio between 80 and 120, aluminum can fully and economically fix nitrogen atoms to form fine and dispersed AlN. The precipitation of AlN in the high-temperature austenite region has the following effects: (1) efficient purification: completely eliminating harmful inclusions such as coarse TiN; (2) effective grain refinement: pinning austenite grain boundaries and inhibiting grain growth; (3) stabilizing and toughening the matrix: moderately increasing the martensitic transformation initiation temperature (Ms) is beneficial to promoting self-tempering and improving the toughness of the matrix.

[0024] Silicon (Si): 0.30-0.50%. Si is an effective solid solution strengthening element that can improve strength and inhibit carbide coarsening. However, excessive Si content (>0.50%) will exacerbate surface red rust defects during hot rolling and affect surface quality.

[0025] Manganese (Mn): 0.40-1.00%. Mn is a key element for improving hardenability, stabilizing austenite, delaying the transformation of pearlite and bainite, and ensuring a fully martensitic structure after hot forming. However, excessive Mn content (>1.0%) will exacerbate center segregation during continuous casting, and worsen bending and welding properties.

[0026] Chromium (Cr): 0.15-0.50%. Cr can improve hardenability, enhance resistance to tempering softening, and improve corrosion resistance. However, Cr is a strong carbide-forming element. Excessive Cr (>0.5%) will combine with C to form coarse carbides, consume effective dissolved carbon, and reduce martensitic hardness.

[0027] Boron (B): 0.001-0.004%. Boron dissolves in high-temperature austenite and segregates towards the austenite grain boundaries, reducing grain boundary energy, inhibiting the heterogeneous nucleation of ferrite and pearlite at grain boundaries, delaying the high-temperature phase transformation process, and significantly improving the hardenability of steel. Too low a B content leads to insufficient hardenability, while too high a B content leads to excessive grain boundary segregation, forming continuous borides, causing hot cracking and a sharp drop in plasticity. B readily combines with N to form BN, reducing the effective solid solubility of B. Therefore, this invention utilizes strong nitride-forming elements such as Al, taking advantage of the much greater affinity of Al for N than that of B, preferentially forming AlN, thus protecting B from being fixed by N. Traditionally, Ti is used, which easily forms large-particle TiN, reducing the toughness of the material.

[0028] Nickel (Ni): ≤0.5%. Ni is an element that can improve strength and plasticity without compromising toughness, and improves cold bending performance. However, Ni is expensive and belongs to precious alloying elements, so it is not advisable to use too much from a cost control perspective.

[0029] Through the above design, this invention not only achieves an optimized match between strength and toughness through fine grain strengthening and precipitation strengthening, but also actively constructs a multi-layered and stable irreversible hydrogen trap network by introducing high-density, uniformly distributed nanoscale microalloyed carbides (NbC and VC), fundamentally enhancing the material's ability to resist hydrogen-induced delayed fracture.

[0030] Compared with the prior art, the beneficial effects of the present invention are as follows: The present invention relates to a 2300MPa grade high-strength and high-toughness hot-formed steel and its preparation method. Based on the distribution of strengthening contribution and the control of graded precipitation, the content of C, Al, V and Nb is determined. The innovative alloying concept, combined with optimized process, starts from the source of composition design to achieve precise and synergistic control of microstructure (especially martensitic matrix, second phase and hydrogen trap). The steel product obtained can simultaneously meet the requirements of ultra-high strength above 2300MPa, good plasticity, excellent cold bending properties, low hydrogen embrittlement sensitivity, and cost-effectiveness.

[0031] The 2300MPa grade high-strength and high-toughness hot-formed steel of this invention exhibits excellent synergistic properties. Its tensile strength is consistently maintained above 2300MPa, its elongation after fracture is ≥5%, its ultimate cold bending angle is ≥45°, and its unique irreversible hydrogen trap network makes its resistance to hydrogen embrittlement far exceed that of similar materials.

[0032] This invention is scientifically designed and cost-effective. It is based on theoretical models and quantitative relationships, and its performance is predictable with a wide process window. Furthermore, it uses economical microalloys (Al, V, Nb) to replace expensive elements, resulting in high cost-effectiveness.

[0033] The present invention has a clean and homogeneous structure, eliminating coarse TiN inclusions. The structure consists of fine lath martensite and dispersed nano-precipitates, resulting in good performance consistency. Attached Figure Description

[0034] Figure 1 This is a schematic diagram of the preparation process of the 2300MPa grade high-strength and high-toughness hot-formed steel of the present invention; Figure 2 This is a schematic diagram of the microalloying graded precipitation mechanism of steel products during hot stamping in a specific embodiment of the present invention; in the figure, A represents austenitic structure and M represents martensite structure; Figure 3 The image shows the fracture morphology of the steel product in Comparative Example 4 of this invention, captured by scanning electron microscopy. In the image, a is a secondary electron image and b is an energy dispersive spectroscopy (EDS) image. Figure 4 This is a measured continuous cooling phase transition (CCT) curve of Embodiment 1 of the present invention; Figure 5 The images show a comparison of the original austenite grains of the steel products in Example 1 and Comparative Example 1; where a is the original austenite grain diagram of the steel product in Example 1, and b is the original austenite grain diagram of the steel product in Comparative Example 1. Figure 6 This is a comparison of the room temperature tensile stress-strain curves of the steel products in Example 1 and Comparative Example 1. Detailed Implementation

[0035] To better explain the present invention, the main contents of the present invention are further illustrated below with reference to the accompanying drawings and specific embodiments, but the contents of the present invention are not limited to the following embodiments.

[0036] This invention discloses a 2300MPa grade high-strength and high-toughness hot-formed steel, comprising the following chemical composition by mass percentage: C: 0.46-0.52%, Si: 0.30-0.50%, Mn: 0.40-1.00%, Cr: 0.15-0.50%, B: 0.001-0.004%, Ti: 0.01-0.03%, Al: 0.15-0.50%, V: 0.10-0.20%, Nb: 0.020-0.050%, Ni≤0.5%, P≤0.015%, S≤0.010%, N≤0.005%, with the balance being Fe and unavoidable impurities; wherein, 4≤V / Nb≤8, 80≤Al / N≤120; V, Nb, Al, and N represent the mass percentages of the corresponding elements; and the C content satisfies: C design (wt.%) = [ target -( Fe + + pre + gb )] / (k c ×ρM ); In the formula, C design The carbon content is designed in wt.%; target The target tensile strength is expressed in MPa. Fe The lattice strength of pure iron is taken as 200 MPa; k1, k2, k3, and k4 are the strengthening coefficients of the substitutional solid solution elements Si, Mn, Cr, and Ni, respectively, taken as 50 MPa / wt.%, 30 MPa / wt.%, 20 MPa / wt.%, and 10 MPa / wt.%; X1, X2, X3, and X4 are the mass percentages (wt.%) of the substitutional solid solution elements Si, Mn, Cr, and Ni. pre The precipitation strengthening contribution is mainly provided by nano-VC, where α is the coefficient of precipitation strengthening contribution, and α is taken as 350 MPa / , pre 350 MPa This represents the mass percentage of V, expressed in wt.%. gb The contribution to grain refinement and strengthening is mainly provided by AlN and Nb(C,N) grain refinement, amounting to 100 MPa; k c The interstitial solid solution strengthening coefficient of carbon is 5000 MPa / wt.%; ρ M The retention coefficient of carbon in the solid solution after the martensitic phase transformation is set to 0.8 because the present invention is designed to capture some carbon atoms by VC precipitation, in order to reduce the effective solid solution carbon concentration.

[0037] like Figure 1 As shown, the preparation method of the 2300MPa grade high-strength and high-toughness hot-formed steel of the present invention includes the following steps: Smelting and Casting: Smelting is carried out in a 100kg vacuum induction furnace according to the target composition and its designed content. The amount of Al added is determined by measuring the N content, ensuring that the final composition meets the requirement of 80 ≤ Al / N ≤ 120. The mixture is then cast into ingots. The designed C content in the target composition satisfies the above formula: C design =[ target -( Fe + + pre + gb )] / (k c ×ρ M ); Hot rolling: The ingot is heated to 1150-1250℃ at a heating rate of 5-10℃ / min and held for 2-4 hours to allow the alloying elements (especially Nb) to fully dissolve. Then, multiple hot rolling passes are performed, with a total reduction rate ≥90% and a final rolling temperature ≥950℃. This stage promotes the precipitation of AlN and Nb(C,N), pins grain boundaries, and inhibits excessively coarse grains. After rolling, the ingot is rapidly cooled to 550-650℃ at a cooling rate of ≥15℃ / s, preferably ≥20℃ / s, more preferably 20-50℃ / s, and held for 1-3 hours. Finally, it is cooled to room temperature in the furnace to obtain a hot-rolled plate with a thickness of 2-5 mm. The room temperature is 0-25℃.

[0038] Cold rolling and annealing: After pickling, the hot-rolled plate is cold-rolled with a reduction rate of 30-60%; then the cold-rolled plate is heated to 600-800℃ at a heating rate of 5-10℃ / s and held for 3-10 minutes for annealing. Afterwards, it is cooled to room temperature in the furnace to eliminate residual stress during the cold rolling process and promote the precipitation of VC.

[0039] Hot forming and quenching: The annealed steel plate is heated to 900-950℃ at a heating rate of not less than 10℃ / s, preferably 10-100℃ / s, and held for 3-5 minutes; after in-mold stamping, it is rapidly quenched with a quenching cooling rate of ≥20℃ / s, in order to obtain a lath martensitic matrix. After cooling to room temperature, a high-strength and high-toughness hot-formed steel of 2300MPa grade is obtained.

[0040] The metallographic structure formation process of this invention is as follows: During hot rolling, the slab is treated at a high temperature of 1200-1250℃ to ensure sufficient dissolution of alloying elements, followed by rolling to refine the austenite grains. Simultaneously, fine AlN and undissolved Nb(C,N) particles precipitate and pin the grain boundaries, resulting in a fine-grained bainite structure with a small amount of ferrite / pearlite. After cold rolling, the hot-rolled plate undergoes elongation and work hardening; subsequent annealing leads to recovery and recrystallization, forming equiaxed ferrite grains, while nanoscale VC particles precipitate, eliminating residual stress and producing precipitation strengthening. Finally, the annealed plate undergoes high-temperature holding to enter the high-temperature austenitization stage, transforming the structure into fine austenite (with partial dissolution of VC); through in-die stamping followed by rapid quenching, the austenite transforms into lath martensite, with microalloying precipitates in stages, ultimately yielding an ultra-high strength structure with a lath martensite matrix and undissolved VC particles.

[0041] Among them, the microalloying graded precipitation mechanism of steel products during hot forming and quenching is described in [reference needed]. Figure 2 .like Figure 2As shown, during the high-temperature austenitization stage (corresponding to the heating and holding process in hot stamping), the fine AlN and undissolved Nb(C,N) particles precipitated during hot rolling effectively suppress the migration of austenite grain boundaries and the growth of recrystallized grains through the Zener pinning effect, achieving grain boundary pinning and grain refinement. At the same time, the high-density grain boundaries and subgrain boundaries introduced by the pinning effect provide more nucleation sites for subsequent phase transformation, promoting the uniformity and refinement of martensite transformation.

[0042] The material then enters the quenching phase transformation stage (corresponding to the rapid cooling process after hot stamping). In the low-temperature region, nanoscale VC phases are dispersed and precipitated within the martensitic laths. These phases significantly improve the matrix strength through precipitation strengthening and also act as irreversible hydrogen traps to enhance the material's resistance to hydrogen embrittlement. Finally, rapid quenching is performed at a cooling rate of at least 20°C / s to obtain a fine lath martensitic matrix, thereby synergistically improving the material's strength, toughness, and service safety.

[0043] The present invention will be further described in detail below through specific embodiments and comparative examples.

[0044] The design yield strength and corresponding design chemical composition and content (unit, wt.%) of the high-strength and high-toughness hot-formed steels in Examples 1-5 are shown in Table 1. The preparation methods are as follows: (1) Smelting: During the smelting process, the alloy composition is sampled and analyzed multiple times until the target value is met before casting into ingots, ensuring 80≤Al / N≤120, 4≤V / Nb≤8, and an ingot thickness of 180mm. The designed carbon content is calculated using the following formula: C design (wt.%) = [ target -( Fe + + pre + gb )] / (k c ×ρ M ) Where, σ Fe =200MPa; k1, k2, k3, and k4 are taken as 50, 30, 20, and 10 MPa / wt.%, respectively. =50×X1+30×X2+20×X3+10×X4, representing the strengthening contribution of the substitution solid solution elements Si, Mn, Cr, and Ni, in MPa; α is the coefficient of precipitation strengthening contribution, α is taken as 350 MPa / , pre =350 MPa: Contribution to enhanced precipitation of nano-VC (based on V content, empirical model). gb =100MPa: Grain refinement strengthening contribution (provided by AlN and Nb(C,N)); k c =5000MPa / wt.%: Interstitial solid solution strengthening coefficient of carbon; ρ M =0.8: Retention coefficient of carbon in solid solution after martensitic phase transformation (some carbon is captured by VC). That is, the formula simplifies to: C design (wt.%) = [ target -(200+ +350 +100)] / (5000×0.8) =[ target -(300+ +350 )] / (4000)。

[0045] (2) Hot rolling: The ingot is heated to 1200℃ at a heating rate of 8℃ / min and held for 3 hours. Then it is hot rolled in multiple passes. The final rolling temperature is 960℃ and the total reduction rate of hot rolling is 98.6%. Then it is cooled to 600℃ at a cooling rate of 25℃ / s and held for 2 hours before being cooled in the furnace to obtain a hot-rolled plate with a final thickness of 2.5mm.

[0046] (3) Cold rolling and annealing: After the hot-rolled plate is pickled to remove the surface oxide scale, it is cold rolled in multiple passes until the thickness is 1.5 mm, and the cold rolling reduction rate is 40%. Then, the temperature is raised to 750°C at 8°C / s and held for 5 minutes for annealing. The steel plate is then cooled in the furnace to obtain the steel plate.

[0047] (4) Quenching: The steel plate is heated to 920°C at a rate of 15°C / s and held for 4 minutes before die quenching. The quenching cooling rate is ≥20°C / s to obtain hot-formed steel products.

[0048] Table 1: Chemical composition and content of the examples, with the balance being Fe and unavoidable impurity elements. Continued from Table 1: Chemical composition and content of the examples, with the balance being Fe and unavoidable impurity elements. The actual chemical composition and content (unit, wt.%) of the steel products in Examples 1-5 and Comparative Examples 1-4 are shown in Table 2. The preparation method of the steel products in the comparative examples is the same as that in the examples. The actual carbon content of the smelting is basically consistent with the theoretical calculation value. The deviations of the actual carbon content of Examples 1-5 (0.48%, 0.49%, 0.52%, 0.46%, 0.48%) from the theoretical carbon content of Table 1 (0.475%, 0.487%, 0.513%, 0.462%, 0.481%) are all within ±0.015 wt.%, with a maximum deviation of only 0.014 wt.%, which fully demonstrates the accuracy and reliability of the carbon content addition of the present invention.

[0049] Table 2: Actual chemical composition and content of the examples and comparative examples, with the balance being Fe and unavoidable impurity elements. Table 2 (continued): Chemical composition and content of the examples and comparative examples, balance being Fe and unavoidable impurity elements. The performance of the steel products in the examples and comparative examples was tested, and the results are shown in Table 3. The test methods for yield strength, tensile strength and elongation after fracture were GB / T 228.1-2021 "Metallic materials - Tensile testing - Part 1: Test method at room temperature"; the test method for cold bending angle was GB / T 232-2024 "Metallic materials - Bending test method"; the constant load delayed fracture test method was GB / T 40403-2021 "Corrosion of metals and alloys - Determination of resistance to stress corrosion cracking of metals by four-point bending method", after being loaded to 100% yield strength, placed in air and the cracking time was recorded.

[0050] Table 3: Performance results of steel products from Examples 1-5 and Comparative Examples 1-4 The target strengths (2400, 2450, 2500, 2350, 2400 MPa) set in Table 1 were compared with the measured tensile strengths (2383, 2427, 2526, 2332, 2363 MPa) in Table 3. The deviations between the measured strengths and the target strengths in each embodiment were -17 MPa, -23 MPa, +26 MPa, -18 MPa, and -37 MPa, respectively, with relative errors all less than 2%. Furthermore, based on the actual tensile strength of each steel product in the examples and the carbon content calculation formula mentioned above, the calculated carbon content is 0.471%, 0.483%, 0.517%, 0.458%, and 0.471%, respectively, which are 0.009%, 0.007%, 0.003%, 0.002%, and 0.009% lower than the actual carbon content, respectively. The relative error is less than 2%, indicating that the steel products prepared based on the composition ratio designed by this formula have a very small error between the actual carbon composition and the C content calculated by the carbon content calculation formula based on the actual tensile strength, further verifying the feasibility of the carbon content formula. This shows that the carbon content design formula based on the strengthening contribution allocation model of this invention has extremely high predictive accuracy and can quantitatively guide the precise addition of carbon content according to the target strength and alloy composition, thereby achieving synergistic control of composition-process-performance, and providing a reliable theoretical basis for the rational design of ultra-high strength hot-formed steel.

[0051] As shown in Table 3, the yield strength of Examples 1-5 reached 1449-1544 MPa, the tensile strength was 2332-2526 MPa, the elongation after fracture was 5.2%-6.4%, and the cold bending angle was 52.3-61.1° perpendicular to the rolling direction and 46.2-50.7° parallel to the rolling direction. The overall performance was significantly better than that of Comparative Examples 1-4. Under 100% yield strength loading, the cracking time of all examples exceeded 300 hours.

[0052] Comparative Example 1 uses a traditional medium-carbon + Mo / Ni system design: the C content is only 0.38%, lacking V / Nb precipitation strengthening, and the Al / N ratio is only 4, resulting in coarse grains reaching 32μm and a tensile strength of only 2171MPa, failing to reach the 2300MPa level. Although expensive Mo and Ni are added, the cost is high but the strength is insufficient.

[0053] Comparative Example 2 adopted a high-C+high-Ti design: the Ti content was as high as 0.052%, and the Al / N ratio was only 13, forming a large number of coarse TiN inclusions, which became crack initiation sites and hydrogen enrichment areas. This resulted in an elongation after fracture of only 4.1%, a cold bending angle of 39.6°, hydrogen embrittlement cracking in only 12 hours, and extremely poor plasticity, bending properties, and resistance to hydrogen embrittlement.

[0054] Comparative Example 3 employs a low V / Nb ratio + high Ti + high Si design: the V / Nb ratio is only 1.8, far lower than the 4-8 of this invention, resulting in insufficient VC precipitation and the presence of coarse TiN inclusions. The tensile strength is 2303 MPa, barely meeting the standard, but the elongation is only 4.5%, the cold bending angle is 43.6°, and the hydrogen embrittlement cracking time is only 27 hours, indicating poor strength-toughness matching.

[0055] Comparative Example 4 employs a high V / Nb ratio + high Al / N ratio design: the V / Nb ratio is 10.5, and the Al / N ratio is 144, exceeding the V / Nb ratio of 4-8 and the Al / N ratio of 80-120 of this invention. The tensile strength is 2341 MPa, reaching the 2300 MPa level, but the elongation after fracture and cold bending angle are poor, only 3.7%, with a cold bending angle of 36.9°. Hydrogen embrittlement cracking occurs in only 49 hours, indicating poor strength-toughness matching. The fracture morphology of the steel ingot from Comparative Example 4 was observed using a scanning electron microscope, and the results are as follows... Figure 3 As shown. By Figure 3 It can be seen that a large number of Al2O3 inclusions were observed in the steel product matrix, resulting in poor plasticity and cold bending performance.

[0056] Therefore, compared with the comparative examples, the present invention avoids a sharp drop in plasticity while maintaining ultra-high strength. The tensile strength of the embodiment is superior to that of comparative example 1, the cold bending formability and directional uniformity are significantly better than those of comparative examples 2-4, and the resistance to delayed fracture is much higher than that of comparative examples 2-4. In summary, the present invention successfully achieves the synergy of ultra-high strength, good plasticity, excellent cold bending performance and outstanding resistance to delayed fracture, solving the technical problem of simultaneously achieving high strength, formability and resistance to hydrogen embrittlement.

[0057] In Example 1 of this invention, during the preparation of the steel product, continuous cooling phase transformation was experimentally measured according to YB / T 5128-2018 "Determination of Continuous Cooling Transformation Curve of Steel - Expansion Method". The results are as follows: Figure 4 As shown. By Figure 4 It is known that in Example 1, Ac3 = 877℃ and Ac1 = 776℃. When the cooling rate is ≥15℃ / s, the diffusion-type transformation is basically suppressed, and austenite directly undergoes martensitic transformation. The martensitic transformation temperature is 352℃. The higher martensitic transformation temperature can promote self-tempering and improve the toughness of martensite. In addition, the martensitic transformation ends at approximately 140℃, and no significant phase transformation volume change occurs during subsequent cooling to room temperature, which is beneficial for maintaining the dimensional accuracy of precision parts. To ensure the stability of the martensitic transformation process, the cooling rate is limited to greater than or equal to 20℃ / s.

[0058] The steel products of Example 1 and Comparative Example 1 were tested for their original austenite grains according to GB / T 6394-2017 "Method for Determination of Average Grain Size of Metals". The results are as follows: Figure 5 As shown. By Figure 5It can be seen that the original austenite grain size in Example 1 is approximately 6 μm, while the original austenite grain size in Comparative Example 1 is approximately 32 μm, showing a significant difference. This indicates that the present invention, by rationally utilizing the grain boundary pinning effect of precipitated phases, can effectively suppress the growth of original austenite grains. High-temperature hot rolling at 1200-1250℃ precipitates AlN and Nb(C,N) to refine the grains, and during quenching and cooling, nanoscale VC phases are dispersed and precipitated within the martensite laths in the low-temperature region, forming a synergistic and complementary effect within the precipitation temperature window, thereby achieving significant grain refinement and improving the strength-toughness balance of the material.

[0059] The steel products of Example 1 and Comparative Example 1 were tested for room temperature tensile stress-strain curves according to GB / T228.1-2021 "Metallic Materials - Tensile Testing - Part 1: Room Temperature Test Method". The results are as follows. Figure 6 As shown. By Figure 6 It can be seen that the yield strength and tensile strength of Example 1 are 1472 MPa and 2383 MPa, respectively, and the elongation after fracture is 5.6%. The yield strength and tensile strength of Comparative Example 1 are 1455 MPa and 2171 MPa, respectively, and the elongation after fracture is 5.2%. The results show that, compared with Comparative Example 1, Example 1 has a slightly improved yield strength, a significantly improved tensile strength of about 200 MPa, and a slightly increased elongation after fracture, showing a better balance between strength and plasticity.

[0060] In summary, all embodiments of the present invention have achieved a stable breakthrough in the synergistic effect of "strength, plasticity, and resistance to hydrogen embrittlement". Specifically, the tensile strength of the material is greater than 2300 MPa, the elongation after fracture is greater than 5%, and the cold bending angle in the parallel rolling direction is ≥45°. In particular, in the constant load delayed fracture test simulating harsh service conditions (applying 100% nominal yield stress), all embodiments did not crack after more than 300 hours. Their resistance to hydrogen embrittlement is far superior to that of the comparative examples, and the design goals have been fully achieved.

[0061] The alloy system developed in this invention is based on economical microalloying elements such as Al, V, and Nb. Its preparation process is highly compatible with mainstream production lines in the existing steel industry, offering a wide production window and stable controllability. This solution significantly improves the overall performance and service safety of materials while effectively controlling production costs. It provides an advanced material solution that combines high performance and cost-effectiveness for the next generation of automotive body structural components aimed at lightweighting and ultra-high safety standards, possessing significant industrial application value.

[0062] This invention represents a systematic innovation in the field of ultra-high strength hot-formed steel. Its core lies in proposing a novel synergistic design concept of "clean nitrogen fixation - hierarchical precipitation - hydrogen trap engineering": In terms of composition, it abandons traditional Ti nitrogen fixation and innovatively adopts a high Al ratio (80≤Al / N≤120) to achieve complete nitrogen fixation and purification, eliminating coarse TiN inclusions and preventing the formation of harmful micron-sized TiN inclusions from the source; in terms of microstructure, it pioneers a "V / Nb spatiotemporal hierarchical precipitation" mechanism (4≤V / Nb≤8), refining Nb(C,N) grains and allowing VC to precipitate nanoscale within martensite, synergistically achieving grain refinement and precipitation strengthening. This allows Nb and V elements to perform their respective functions at high and low temperatures, respectively undertaking grain refinement strengthening and nanoscale precipitation strengthening, and both being converted into effective hydrogen traps; in terms of performance, the aforementioned nanoscale precipitated phases are designed as irreversible hydrogen traps, constructing a highly efficient hydrogen capture network. Ultimately, with a tensile strength > 2300 MPa, it achieved for the first time high elongation after fracture, excellent bending properties, and outstanding resistance to hydrogen embrittlement, breaking through the inherent contradiction between strength, toughness, and resistance to hydrogen embrittlement.

[0063] The above description is merely a preferred embodiment of the present invention and does not limit the patent scope of the present invention. All equivalent structural transformations made based on the inventive concept of the present invention and the description and drawings, or direct / indirect applications in other related technical fields, are included within the patent protection scope of the present invention. Any other parts not described are considered prior art.

Claims

1. A 2300MPa grade high-strength and high-toughness hot-formed steel, characterized in that, The chemical composition includes the following components by mass percentage: C: 0.46-0.52%, Si: 0.30-0.50%, Mn: 0.40-1.00%, Cr: 0.15-0.50%, B: 0.001-0.004%, Al: 0.15-0.50%, V: 0.10-0.20%, Nb: 0.020-0.050%, N≤0.005%; where: 4≤V / Nb≤8, 80≤Al / N≤120, and V, Nb, Al, and N represent the mass percentage of the corresponding elements.

2. The 2300MPa grade high-strength and high-toughness hot-formed steel according to claim 1, characterized in that, In the high-strength and high-toughness hot-formed steel, the C content satisfies the following formula: C design =[ target -( Fe + + pre + gb )] / (k c ×ρ M ); In the formula, C design The carbon content is designed in wt.%; target The target tensile strength is expressed in MPa. Fe ρ represents the lattice strength of pure iron, in MPa; k1, k2, k3, and k4 are the strengthening coefficients of the substitutional solid solution elements Si, Mn, Cr, and Ni, respectively, in MPa / wt.%; X1, X2, X3, and X4 are the mass percentages of the substitutional solid solution elements Si, Mn, Cr, and Ni, respectively, in wt.%. pre The precipitation enhancement contribution is expressed in MPa. gb Contribution to grain refinement strengthening, in MPa; k c ρ is the interstitial solid solution strengthening coefficient of carbon, expressed in MPa / wt.%; M is the retention coefficient of carbon in the solid solution after the martensitic phase transformation, which is dimensionless.

3. The 2300MPa grade high-strength and high-toughness hot-formed steel according to claim 2, characterized in that, The Fe The value range is 180-220 MPa; the value ranges of k1, k2, k3, and k4 are 45-55, 28-32, 18-22, and 8-12 MPa / wt.%, respectively. pre =α MPa, where α is the coefficient of precipitation strengthening contribution, and the value of α ranges from 330 to 370 MPa / , This represents the mass percentage of V, in wt.%; gb The value range is 80-120 MPa; k c The value range is 4800-5200 MPa / wt.%; ρ M The value range is 0.75-0.

85.

4. The 2300MPa grade high-strength and high-toughness hot-formed steel according to claim 1, characterized in that, The chemical composition of the high-strength and high-toughness hot-formed steel also includes: Ti: 0.01-0.03%, Ni≤0.5%, P≤0.015%, S≤0.010%, with the balance being Fe and unavoidable impurities.

5. The 2300MPa grade high-strength and high-toughness hot-formed steel according to any one of claims 1-4, characterized in that, The high-strength and high-toughness hot-formed steel has a tensile strength ≥2300MPa, an elongation after fracture ≥5%, and an ultimate cold bending angle ≥45°.

6. A method for preparing the 2300MPa grade high-strength and high-toughness hot-formed steel according to any one of claims 1-5, characterized in that, include: The high-strength and tough hot-formed steel is obtained by smelting and casting, hot rolling, cold rolling and annealing, hot forming and quenching. The smelting process ensures that the final composition meets the requirement of 80 ≤ Al / N ≤ 120.

7. The preparation method according to claim 6, characterized in that, During the hot rolling process, the cast ingot obtained from casting is first heated to 1150-1250℃ and held for 2-4 hours; then hot rolling is carried out with a total reduction rate of ≥90% and a final rolling temperature of ≥950℃; after hot rolling, it is rapidly cooled to 550-650℃ at a cooling rate of ≥15℃ / s and held for 1-3 hours, and finally cooled in the furnace to obtain the hot-rolled plate.

8. The preparation method according to claim 6, characterized in that, During the cold rolling process, the reduction rate is 30-60%; during the annealing process, the cold-rolled sheet obtained by cold rolling is first heated to 600-800℃ at a heating rate of 5-10℃ / s, held for 3-10 minutes, and then cooled in the furnace to obtain the steel sheet.

9. The preparation method according to claim 6, characterized in that, In the thermoforming process, the annealed steel plate is heated to 900-950°C at a heating rate of 10°C / s or higher, held at that temperature for 3-5 minutes, and then thermoformed by in-mold stamping.

10. The preparation method according to any one of claims 6-9, characterized in that, The quenching process includes directly quenching and cooling the hot-formed steel plate to obtain the high-strength and high-toughness hot-formed steel; the cooling rate of the quenching and cooling process is ≥20℃ / s.

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