Low-carbon high-toughness hot stamping forming component and steel sheet

Abstract

The present invention relates to a high-toughness hot stamping component and its steel sheet with low carbon content. The steel sheet matrix of the hot stamping component comprises, by mass percentage:Carbon (C): 0.053% ≤ C ≤ 0.10%, Silicon (Si): 0.05% ≤ Si ≤ 0.30%, Manganese (Mn): 1.81% ≤ Mn ≤ 2.7%, Chromium (Cr): 0.01% ≤ Cr ≤ 0.7%, Aluminum (Al): 0.01% ≤ Al ≤ 0.5%, Boron (B): 0.0005% ≤ B ≤ 0.005%, Titanium (Ti): 0.015% ≤ Ti ≤ 0.05%, Niobium + Vanadium (Nb + V): 0 ≤ Nb + V ≤ 0.2%, Phosphorus (P): 0.001% ≤ P ≤ 0.100%, Sulfur (S): 0.0001% ≤ S ≤ 0.100%, Iron (Fe): ≥ 95% and inevitable impurities with the condition that Mn + 0.26Si + 1.3Cr ≥ 2.20%.The microstructure of the steel sheet matrix, by area percentage, consists of: bainite: < 5%, austenite: < 3%, ferrite: < 3%, Nb-V-Ti carbides: < 0.2%, the remainder is martensite.

Description

Technical Field

[0001] The present invention relates to a low-carbon and high-toughness hot stamping forming component and its steel sheet.Background

[0002] As the highest strength-grade material in automotive steels, the application proportion of hot stamping steels in automotive materials is increasing annually. Correspondingly, the automotive industry has higher requirements for the strength and toughness of hot stamping steels. Room-temperature tensile tests (GB / T 228.1 standard) are commonly used to measure material strength, with yield strength and tensile strength reflecting resistance to tensile deformation failure. Static three-point bending tests (VDA 238-100 standard) are widely employed to assess material toughness by measuring fracture strain.

[0003] To meet automotive safety requirements, specific vehicle regions (e.g. the lower end of B-pillars) must absorb significant impact energy during collisions without brittle fracture. These components also require high yield strength and work-hardening capacity to prevent premature deformation. While 22MnB5 steel is widely used for its>1400MPa tensile strength, its low toughness fails to meet stringent collision regulations and energy-absorption demands, posing safety risks.

[0004] CN107810281A discloses a press-hardened steel and components manufactured from such steel, primarily intended for structural elements in automotive applications serving intrusion prevention or energy absorption functions. The steel achieves satisfactory tensile strength, bendability, and weld zone toughness by maintaining carbon content within 0.062% ≤C≤0.095%. Hardenability elements (C, Mn, Si, Cr) are controlled to meet 1.5%≤ (C+Mn+Si+Cr)≤2.7%, ensuring a microstructural balance between martensite and ductile bainite in hot stamped parts. The patent emphasizes increased bainite content to enhance component toughness and bendability, explicitly requiring Cr≤0.1% since chromium addition hinders bainite formation. However, two critical challenges exist: 1) Current hot stamping processes struggle to achieve the required bainite-martensite dual-phase microstructure; 2) Excessive bainite reduces yield strength, risking premature deformation failure during collisions and compromising vehicle safety.

[0005] CN104769138A proposes that forming a decarburized zone with a depth p50% (defined as the depth where the carbon content reaches 50% of the substrate steel's carbon content) ranging from 6-30µm on the substrate steel surface prior to coating significantly improves the bendability of final hot-formed components. For instance, this method enables a VDA bending angle exceeding 55°for 1.8 mm-thick galvanized 22MnB5 hot stamped steel sheets. However, the critical decarburization depth is highly dependent on annealing conditions before coating, increasing production complexity. Additionally, the patent overlooks adverse effects of the decarburized layer, including reduced material strength, compromised surface hardness, elevated peak load during bending, and potential quality issues in subsequent coating processes.

[0006] CN1717499A discloses a high-strength steel plate for cold working forming and its manufacturing method. To ensure excellent workability and high strength, the steel composition is controlled to C: 0.05-0.15%, Si: 0.3-2.0%, Mn: 1.0-2.8%, Al: 0.005-0.5%, N:≤0.006%, with the remainder being Fe and inevitable impurities, while maintaining Mn / C≥12 and Si / C≥4. The ratio design of C, Mn, and Si aims to retard pearlite and cementite phase transformations during post-coating alloying treatment. The steel's microstructure predominantly features ferrite with 3-20% martensite and retained austenite, achieving a tensile strength of 500-900 MPa and elongation of 27-38%, balancing high strength and workability. However, during hot stamping forming (typically involving 870-930°C holding followed by die quenching), the desired microstructure cannot be stabilized, preventing simultaneous attainment of high strength and workability. Furthermore, the Mn / C≥12 and Si / C≥4 ratios, optimized for alloying processes (450-520°C), become ineffective in hot stamping forming (which excludes alloying). Consequently, CN1717499A's high-strength steel plate is unsuitable for producing high-strength, high-toughness hot stamped components.

[0007] In light of the above, there remains a critical need to develop a hot stamping steel, hot stamped components, and corresponding forming methods that simultaneously achieve high strength and high toughness to enhance automotive safety.Summary of the Invention

[0008] This invention is made in view of the aforementioned problems in the prior art.

[0009] One objective of the present invention is to provide a hot stamping steel sheet with improved hardenability under low carbon content. The steel matrix of the hot stamping steel sheet comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, Fe≥95% and the remainder being inevitable impurities. The Mn, Si, and Cr contents satisfy the condition: Mn+0.26Si+1.3Cr≥2.20%.

[0010] Preferably, the steel matrix further optionally contains at least one of the following elements by mass percentage: 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001% ≤Zr≤0.01%, 0.0001%≤REM≤0.01%.

[0011] Further preferably, 0.0001%≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30%.

[0012] The steel matrix of the hot stamping forming steel sheet according to the present invention comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤ Si≤0.30%, 1.81%≤Mn≤ 2.7%, 0.01%≤Cr≤0.7%, 0.01%≤ Al≤0.5%, 0.0005%≤ B≤ 0.005%, 0.015%≤ Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, with the remainder being Fe and inevitable impurities. The Mn, Si, and Cr contents satisfy the condition:Mn+0.26Si+1.3Cr≥2.20%.

[0013] Preferably, the steel matrix of the hot stamping forming steel sheet according to the present invention comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, and at least one of the following elements within the specified ranges: 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01%, 0.0001% ≤REM≤0.01%, with the total content of these optional elements satisfying: 0.0001%≤(W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM)≤0.30%, and the remainder being Fe and inevitable impurities. The Mn, Si, and Cr contents must satisfy:Mn+0.26Si+1.3Cr≥2.20%.

[0014] Under the premise of low carbon content, this invention achieves improved hardenability through synergistic ratio design of Mn, Cr, Si, and auxiliary microalloying elements.

[0015] Preferably, the Mn+0.26Si+1.3Cr ≥2.25%, or even ≥2.50%.

[0016] Preferably, 0.11%≤Cr≤0.5%; and more preferably0.12%≤Cr≤0.4%.

[0017] Preferably, 2.01%≤Mn≤2.50%.

[0018] Preferably, 0.01≤Nb+V≤0.2%.

[0019] Preferably, 0.055%≤C≤0.09% and0.10%≤Si≤0.26%.

[0020] Preferably, Si+Cr≤0.50%, and more preferably Si+Cr≤0.47%, to further improve the coating surface quality of galvanized steel sheets while ensuring more uniform microstructure and superior mechanical properties after hot stamping.

[0021] In this invention, the martensite start temperature (Ms) of the steel sheet is calculated using the alloy element mass percentages according to the formula: Ms = 520 − 320 C − 50 Mn − 5 Si − 20 Ni − 30 Cr − 20 Mo − 5 Cu Preferably, the Ms temperature is controlled within 350-410°C, and more preferably 360-400°.This formula is derived from A.V. Sverdlin and A.R. Ness "Steel Heat Treatment Handbook" (Edited by G.E. Totten and M.A.H. Howes, Marcel Dekker Inc., New York, 1997).

[0022] In this invention, the ferrite-to-austenite transformation finish temperature (Ac 3 ) of the steel sheet is calculated using the alloy element mass percentages according to the formula: Ac 3 = 912 − 250 C − 16 Mn + 48 Si − 2 Cr − 16 Ni + 95 V + 96 Ti + 210 Al − 10 Cu Preferably, the Ac 3 temperature is controlled to ≤900°C. This formula is derived from Zeng Qiyin's study, 〈〈Quantitative Analysis of Alloying Elements' Effects on Phase Transition Points Ac1 and Ac3 in Steels〉〉 (Physical Testing Volume, 1985, 18(5): 47-49) .

[0023] The hot stamping steel sheet may be coated with a metallic coating. Preferably, the metallic coating can be an aluminum-based alloy coating or a zinc-based alloy coating.

[0024] Another objective of the present invention is to provide a hot stamping formed component with high strength and high toughness. The steel sheet substrate of the hot stamping formed component comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, Fe≥95% and inevitable impurities. The Mn, Si, and Cr contents satisfy: Mn+0.26Si+1.3Cr≥2.20%.Microstructurally, the steel sheet substrate contains (by area percentage):<5% bainite, <3% retained austenite, <3% ferrite, <0.2% Nb -V-Ti microalloyed carbides, with the remainder being martensite.

[0025] Preferably, the steel sheet substrate of the hot stamping formed component further optionally contains at least one of the following elements (by mass percentage): 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤ 0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01%, 0.0001%≤REM (Rare Earth Metals)≤ 0.01%.

[0026] Preferably, 0.0001%≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30%.

[0027] Preferably, the steel sheet substrate of the hot stamping formed component comprises, by mass percentage: 0.053%≤C≤ 0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01% ≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤ 0.100%, 0.0001% ≤S≤0.100%, and the remainder being Fe and inevitable impurities. The Mn, Si, and Cr contents satisfy: Mn+0.26Si+1.3Cr≥2.20%.

[0028] Preferably, the steel sheet substrate of the hot stamping formed component comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤ Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, and at least one of the following elements (by mass percentage): 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01%, 0.0001%≤REM (Rare Earth Metals)≤0.01%, with the total content of these optional elements satisfying: 0.0001%≤(W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM)≤0.30%, and the remainder being Fe and inevitable impurities. The Mn, Si, and Cr contents satisfy: M +0.26Si+1.3Cr≥2.20%.

[0029] Compared with existing technologies, this invention optimizes the hot stamping formed component through a synergistic design of Mn, Cr, Si, and other alloying elements under low carbon content, achieving high strength while maintaining high toughness. This meets the dual requirements of energy absorption in specific vehicle body regions and lightweight design.

[0030] Preferably, the combined content of bainite, ferrite, and retained austenite in the steel sheet substrate does not exceed 3%.

[0031] The steel component may be coated with a metallic coating. Preferably, the metallic coating may be an aluminum-based alloy coating or a zinc-based alloy coating.

[0032] The component exhibits a yield strength (YS) of 880-1100 MPa, a tensile strength (TS)≥1000 MPa, elongation≥5%, and fracture strain (ε)≥0.6.

[0033] The component exhibits a yield strength (YS) of 930-1050 MPa, a tensile strength (TS)≥1060 MPa, elongation≥7%, and fracture strain (ε)≥0.7.

[0034] Preferably, the bending energy absorption (W) of the component satisfies: W / t2≥2.88 × 104N / mm, where t is the sheet thickness.

[0035] The present invention further provides a hot stamping method for producing the aforementioned hot stamped formed component, comprising: A) Providing a billet or preformed blank made from the hot stamping steel described above; B) Full austenitization treatment: Heating the billet or preformed blank from step A to at least 900°C and holding at this temperature; C) Hot stamping and forming: After step B, transferring the heated billet, preform to a press for hot deformation and cooling in the mold to<300°C to obtain the hot stamped component. During cooling, the average cooling rate in non-deformed regions(from 700°C to Ms temperature) is≥40°C / s, while deformed regions are cooled at higher average rates.

[0036] Preferably, the present invention further optimizes the hot stamping method by specifying Ms temperature range 350-410°C, Step B parameters: 1) Heating furnace temperature:900-980°C. 2) Holding time (minutes): t×3+10 (where t = billet / preform thickness in mm).

[0037] The present invention further optimizes the hot stamping method by specifying: In step C: 1) Local deformation constraint: At least one region of the billet / preform undergoes ≤10% plastic strain during hot forming. 2) Enhanced cooling rate: The same region is cooled from 700°C to Ms temperature at an average cooling rate ≥50°C / s.

[0038] Further preferably, the present invention optimizes the hot stamping method by specifying: In step C: 1) Enhanced deformation constraint: At least one region of the billet / preform undergoes≤20% plastic strain during hot forming. 2) Accelerated cooling rate: The same region is cooled from 700°C to Ms temperature at an average cooling rate≥60°C / s.

[0039] It is well known that carbon (C), as a common hardenability element in steel, significantly increases strength but negatively impacts toughness. Therefore, reducing carbon content is the most effective way to achieve sufficient toughness. However, low-carbon designs compromise strength, conflicting with vehicle lightweighting requirements. To balance strength and toughness, this invention proposes achieving high strength through phase transformation strengthening, i.e., obtaining a martensitic structure (ideally full martensitic structure with total content exceeding 97%) via hot stamping deformation and mold cooling. Under low-carbon design constraints, optimizing alloy element ratios is critical to ensure sufficient hardenability.

[0040] It is widely recognized that bainite possesses higher toughness than martensite. However, this invention reveals that in low-carbon steels, the solid solution strengthening effect of carbon is weakened, and the formation of bainite involves significant carbide precipitation, further reducing the solid solution strengthening effect in the matrix, thereby compromising material strength. Therefore, this invention aims to avoid bainitic transformation under low-carbon conditions to ensure high strength. Notably, bainite formation is a diffusion-controlled phase transformation, and chromium (Cr), as a strong carbide-forming element, reduces carbon diffusion capacity, effectively suppressing bainite nucleation. This mechanism differs from conventional alloying elements (e.g., Mn) that enhance hardenability by stabilizing austenite thermodynamically. Additionally, silicon (Si), as a solid solution strengthening element, promotes ferrite nucleation and growth, expanding the temperature range for ferrite transformation during cooling. Ferrite formation, however, undermines high strength. To address this, the invention employs a dual strategy under low-carbon design.Control Cr and Si contents:

[0041] 1) Suppress bainite nucleation via Cr's carbide-forming ability. 2) Reduce ferrite promotion by optimizing Si content. Optimize hardenability through the Mn+0.26Si+1.3Cr ratio, ensuring synergistic alloying effects. This approach achieves a balanced ratio of Mn, Si, and Cr, maximizing martensitic content while minimizing non-martensitic phases (ferrite, bainite, retained austenite). The result is a material with high yield strength and tensile strength, coupled with sufficient ductility and toughness, tailored for hot stamping forming components.

[0042] Furthermore, a lower Ac3 temperature(austenitizing temperature) facilitates complete austenitization of the steel plate during hot stamping. By controlling the Ac3 temperature, the final microstructure's ferrite content can be further minimized.

[0043] Furthermore, an optimal Ms temperature ensures the final martensitic matrix achieves both high strength and toughness. While it is widely believed that increasing Ms temperature promotes martensite self-tempering and improves toughness, this invention reveals that in low-carbon steels, the martensitic matrix already exhibits sufficient ductility. However, if Ms exceeds 410°C, excessive self-tempering occurs, leading to reduced strength and toughness. Therefore, under low-carbon conditions, this invention preferentially controls Ms temperature within 350-410°C to balance self-tempering effects, ensuring the martensitic matrix retains high strength and toughness.

[0044] Compared to conventional steels, the hot-stamping formed steel sheet of this invention exhibits improved hardenability. The resulting hot-stamped components achieve: Yield strength (YS): 880-1100 MPa Tensile strength (TS): ≥1000 MPa Elongation: ≥5% Fracture strain (ε): ≥0.6 Bending energy absorption (W / t 2< ): ≥2.88×10 4< N / mm The high yield and tensile strengths ensure lightweight performance while preventing premature deformation failure during collisions. The combination of high strength and excellent ductility-toughness enables effective energy absorption during impact events.Brief Description of the Drawings

[0045] Figure 1: Typical microstructure of the T1 composition hot-stamped component after hot stamping forming. Figure 2: Typical microstructure of the T2 composition hot-stamped component after hot stamping forming. Figure 3: Typical microstructure of the CT1 composition hot-stamped component after hot stamping forming.

[0046] In all figures, M denotes martensite, and B denotes bainite.Detailed Description of the Invention

[0047] The invention will now be described in greater detail below with reference to exemplary embodiments. In this specification, percentages of chemical elements (denoted by %) refer to weight percentages, unless otherwise specified. All preferred embodiments may be freely combined as needed, unless explicitly restricted. All ranges disclosed herein include their endpoints unless expressly stated otherwise. Those skilled in the art will understand that the data and parameters described in the embodiments are exemplary only and do not constitute limitations of the invention.

[0048] The steel substrate of the hot-stamped component according to the present invention comprises the following components by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001%≤S≤0.100%, Fe≥95% and inevitable impurities. The combined content of Mn, Si, and Cr must satisfy:Mn+0.26×Si+1.3×Cr≥2.20%. Microstructure by area percentage: bainite<5%, austenite <3%, ferrite<3%, Nb-V-Ti microalloy carbides<0.2%, and the remainder being martensite.

[0049] Preferably, the steel substrate further comprises at least one of the following elements by mass percentage: 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01%, 0.0001%≤REM≤0.01%.

[0050] Further preferably, 0.0001%≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30%.

[0051] Preferably, the remainder of the steel substrate consists of iron (Fe) and inevitable impurities.

[0052] Preferably, Mn+0.26Si+1.3Cr≥2.25%, and further Mn+0.26Si+1.3Cr≥2.50%.Preferably, 0.11%≤Cr≤0.5%, further preferably, 0.12%≤Cr≤0.4% Preferably, 0.01≤Nb+V≤0.2% Preferably, 0.055%≤C≤0.09%; 0.10%≤Si≤0.26%.Preferably, Si+Cr≤0.50%.Preferably, 1.90%≤Mn≤2.50%, and further2.01%≤Mn≤2.50% Preferably, 0.03%≤Ti≤0.046%.

[0053] Preferably, the martensite start temperature (Ms)of the component is calculated using the alloy element composition by mass percentage according to the formula: Ms=520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu, and must satisfy:350°C≤Ms≤410°C.The austenite finish temperature (Ac 3 ) is calculated as:Ac3=912-250C-16Mn+48Si-2Cr-16Ni+95V+96Ti+210Al-10Cu, and must satisfy: Ac3≤900°C.

[0054] The chemical composition of the hot stamped component according to the present invention is described in detail as follows: C: 0.053~0.10%C: 0.053~0.10%

[0055] Carbon (C) is the most common alloying element in steel for improving strength. Higher carbon content enhances hardenability and strength but reduces fracture strain, thereby decreasing toughness. Additionally, carbon significantly impacts phase transformation characteristics: lower carbon content increases the Ac3 and Ms temperatures. As the foundational design of this invention, the carbon content is controlled to≤0.10% to ensure the hot-stamped component achieves the desired toughness. However, since carbon improves work hardening capacity during deformation, a minimum carbon content is critical. Based on these considerations, the carbon content is optimized within0.053-0.10%, preferably0.055-0.09%.

[0056] Si: 0.05~0.30%, Mn:1.81~2.70%, Cr: 0.01~0.7% and Mn+0.26Si+1.3Cr≥2.20% Under low-carbon design, to ensure maximum martensite generation during hot stamping for high strength, this invention requires the optimization of ratios of cost-effective elements such as Si, Mn, and Cr that enhance hardenability.

[0057] Silicon (Si) dissolves in the steel matrix, enhancing strength and improving hardenability to a certain extent. Therefore, this invention requires Si content ≥0.05%. However, Si content exceeding 0.30% causes severe surface oxidation and decarburization during production, degrading surface quality. Excessive Si also stabilizes ferrite, reducing martensite formation during hot stamping and weakening strength. Thus, the invention specifies 0.05-0.30% Si, preferably 0.10-0.26% Si.

[0058] Manganese (Mn) enhances steel hardenability, increases austenite stability, expands the austenitizing phase field, and lowers the Ac3 temperature. When Mn content is below 1.81%, it cannot offset the elevated Ac3 and Ms temperatures caused by the low carbon content in this invention. This results in incomplete austenitization during the hot stamping heating process and excessive ferrite / bainite formation during cooling, thereby reducing component strength and toughness. Conversely, Mn content exceeding 2.70% may cause segregation during production, negatively impacting ductility and toughness. Therefore, this invention specifies Mn content within 1.81-2.70%, preferably 2.01-2.50%.

[0059] Chromium (Cr) not only enhances steel hardenability but also significantly improves oxidation resistance and prevents surface decarburization. Additionally, the inventors found that Cr, as a strong carbide-forming element, reduces carbon diffusion capacity. Since bainite transformation is a diffusion-controlled phase transition, adding ≥0.01% Cr helps suppress bainite formation, which is complementary to the hardenability enhancement provided by other alloying elements (e.g., Mn). However, Cr content exceeding 0.7% causes severe surface oxidation during hot rolling, compromising surface quality after pickling. Therefore, this invention specifies Cr content within 0.01-0.7%, preferably 0.11-0.5%, and more preferably 0.12-0.40%.

[0060] When the Si, Mn, and Cr contents fall within the aforementioned ranges, if Mn+0.26Si+1.3Cr < 2.20%, insufficient hardenability will result in inadequate strength. Therefore, this invention requires that Mn+0.26Si+1.3Cr≥2.20% when Si, Mn, and Cr are within their specified ranges. Preferably, this sum should be≥2.25%, and most preferably≥2.50%.

[0061] On one hand, since Si and Cr are more prone to oxidation than Fe, excessive addition can cause severe surface oxidation of hot-rolled coils, reducing pickling effectiveness and compromising the coating quality of coated steel sheets. This also affects the interdiffusion between Fe and Al elements at the coating / substrate interface during hot-stamping heating. On the other hand, an optimal Si + Cr ratio suppresses the formation of non-martensitic phases (e.g., ferrite and bainite) after hot stamping, thereby maximizing martensitic structure and achieving superior comprehensive performance. Therefore, this invention preferably requires Si + Cr ≤0.50%, and more preferably ≤0.47%.Al: 0.01~0.5%

[0062] Aluminum (Al) is a strong deoxidizer, typically added in amounts≥0.01% during steel melting to remove oxygen and reduce inclusions. However, excessive Al increases the Ac3 temperature, preventing full austenitization under the heating conditions of this invention's hot-stamping process. Additionally, high Al content exacerbates resistance at the mold nozzle during continuous casting. Therefore, this invention limits Al content to≤0.5%.B: 0.0005~0.005%

[0063] Boron (B) segregates at austenite grain boundaries during hot stamping, inhibiting ferrite formation and significantly enhancing hardenability. Therefore, a certain amount of B is added to components in this invention. However, excessive B causes boron embrittlement, so the steel's B content is controlled within 0.0005-0.005%.Ti: 0.015~0.05%

[0064] Titanium (Ti) forms a strong bond with nitrogen (N). Adding ≥0.015% Ti to hot-stamped steel ensures Ti reacts with dissolved nitrogen to form TiN, preventing N from binding with boron (B) and reducing B's effectiveness. Additionally, Ti carbonitrides enhance strength and refine grain structure, improving toughness. However, excessive Ti (>0.05%) degrades ductility. Thus, Ti content is controlled within 0.015-0.05% in this invention.Nb+V: 0~0.2%

[0065] The addition of niobium (Nb) and vanadium (V) in steel forms carbo-nitrides, which enhance material performance through precipitation strengthening and grain refinement. Therefore, these elements may be added in appropriate amounts in the components of this invention. However, excessive addition of Nb and V increases production costs. Thus, this invention controls the combined content of Nb and V within 0-0.2%, preferably 0.01-0.2%.P: 0.001%~0.100%

[0066] Phosphorus (P) is an unavoidable element in steel. On one hand, P acts as a solid solution strengthening element, relatively cost-effectively enhancing the strength of the steel plate. On the other hand, when P content exceeds 0.100%, it segregates at grain boundaries, leading to significant reductions in toughness and other adverse effects. Therefore, the upper limit of P content is controlled to ≤0.100%, preferably ≤0.050%. Additionally, since reducing P content below 0.001% would increase smelting costs, the lower limit is set to ≥0.001%, preferably ≥0.004%.S: 0.0001%~0.100%

[0067] Similar to phosphorus (P), sulfur (S) is an unavoidable element in steel, forming inclusions such as MnS through reaction with manganese (Mn). When S content exceeds 0.100%, excessive MnS inclusions significantly degrade ductility and toughness, worsening processability. Therefore, the upper limit of S content is controlled to ≤0.100%, preferably ≤0.015%. Additionally, reducing S content below 0.0001% increases smelting costs, so the lower limit is set to ≥0.0001%, preferably ≥0.0005%, and more preferably ≥0.001%.

[0068] In addition to the core elements (C, Si, Mn, Cr, Al, B, Ti, Nb, V, P, S, etc.), the steel substrate of hot-stamped components may also contain the following elements, with their roles, content ranges, and design considerations outlined below.W, Mo, Ni, Cu, Co: 0.01~0.30%

[0069] The addition of tungsten (W), molybdenum (Mo), nickel (Ni), copper (Cu), and cobalt (Co) enhances the hardenability of steel. However, these alloying elements increase production costs, so they should be added in moderation. Additionally, Ni, Cu, and Co improve material toughness. When the content of each element is ≥0.01%, these beneficial effects become evident. Therefore, the lower limits for these five elements are preferably set to ≥0.01% each. To balance cost efficiency and processability, their upper limits are preferably capped at ≤0.30% each. This approach ensures improved toughness while minimizing impacts on hardenability and maintaining machinability.Sn: 0.005~0.300%, Sb: 0.005~0.100%

[0070] When the contents of tin (Sn) and antimony (Sb) are ≥0.005%, they can improve the wettability of coatings. Therefore, the lower limits for Sn and Sb are preferably set to ≥0.005% each. However, excessive Sn (>0.300%) and / or Sb (>0.100%) degrade material ductility. Thus, the upper limits are preferably capped at Sn ≤0.300% and Sb ≤0.100%.

[0071] Ca: 0.0001 ~0.01 % , Mg: 0.0001 ~0.01 % , Zr: 0.0001 ~0.01 % , REM: 0.0001 ~ 0.01%

[0072] Rare Earth Metals (REM), Calcium (Ca), Magnesium (Mg), and Zirconium (Zr) play critical roles in steel performance when their contents are controlled within specific ranges: 1. Minimum Content Requirements REM (≥0.0001%) : Purifies the steel melt by removing inclusions. Modifies impurity morphology and distribution, reducing brittle phases. Refines grain structure, enhancing strength and toughness. Ca / Mg / Zr (≥0.0001%) : Refine inclusions (e.g., sulfides, oxides) via nucleation effects, improving material cleanliness and ductility. 2. Maximum Content Limits REM (≤0.01%): Excess REM (>0.01%) causes over-refinement (grain coarsening) and embrittlement due to intermetallic precipitation. Ca / Mg / Zr (≤0.01%): Higher amounts lead to saturation of effects (no further improvement) and may form harmful precipitates (e.g., CaO, MgO).

[0073] Preferably, to account for alloy cost considerations, the combined content of W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr, and Rare Earth Metals (REM) is controlled within 0.0001-0.30%. Overall, the steel substrate of this invention may contain elements such as W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr, and REM. The presence of these elements does not interfere with the technical solution of the present invention.

[0074] In this embodiment, there are no specific restrictions on other components of the steel substrate for the hot-stamped component. For example, elements such as arsenic (As) may occasionally be introduced from scrap materials, but if within typical ranges, they will not affect the characteristics of the hot-stamped component.

[0075] Preferably, the first exemplary steel substrate for the hot-stamped component comprises the following composition by mass percentage: C: 0.059-0.089%, Si: 0.10-0.30%, Mn: 1.88-2.47%, Cr: 0.08-0.31%, Al: 0.015-0.05%, B: 0.0015-0.0035%, Ti: 0.018-0.046%, Nb + V: 0-0.15%, P: 0.001-0.100%, S: 0.0001-0.100%, Fe ≥95%, and inevitable impurities. Wherein, the Mn, Si, and Cr contents satisfy: 2.20%≤Mn+0.26Si+1.3Cr≤2.82%. Preferably, the first exemplary steel substrate further includes at least one of the following elements by mass percentage: W: 0.01-0.30%, Mo: 0.01-0.30%, Ni: 0.01-0.30%, Cu: 0.01-0.30%, Co: 0.01-0.30%, Sn: 0.005-0.30%, Sb: 0.005-0.100%, Ca: 0.0001-0.01%, Mg: 0.0001-0.01%, Zr: 0.0001-0.01%, REM: 0.0001 -0.01%. More preferably, the total content of W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM is controlled within: 0.0001%≤Total≤0.30%.

[0076] More preferably, the first exemplary steel substrate for the hot-stamped component consists of iron (Fe) and inevitable impurities for the remainder of its composition by mass percentage.

[0077] More preferably, the second exemplary steel substrate for the hot-stamped component comprises the following composition by mass percentage: C: 0.059-0.089%, Si: 0.10-0.26%, Mn: 1.88-2.40%, Cr: 0.11-0.31%, Al: 0.015-0.05%, B: 0.0015-0.0035%, Ti: 0.018-0.046%, Nb + V: 0-0.15%, P: 0.001-0.100%, S: 0.0001-0.100%, Fe ≥95%, and inevitable impurities. Wherein, the Mn, Si, and Cr contents satisfy:2.27%≤Mn+0.26Si+1.3Cr≤2.60%. Preferably, the second exemplary steel substrate further includes at least one of the following elements by mass percentage: W: 0.01-0.30%, Mo: 0.01-0.30%, Ni: 0.01-0.30%, Cu: 0.01-0.30%, Co: 0.01-0.30%, Sn: 0.005-0.30%, Sb: 0.005-0.100%, Ca: 0.0001-0.01%, Mg: 0.0001-0.01%, Zr: 0.0001-0.01%, REM: 0.0001-0.01%. More preferably, the total content of W+Mo+Ni+Cu+Co+Sn+Sb+Ca+ Mg+Zr+REM is controlled within: 0.0001%≤Total≤0.30%.

[0078] More preferably, the second exemplary steel substrate for the hot-stamped component consists of iron (Fe) and inevitable impurities for the remainder of its composition by mass percentage.

[0079] The present invention achieves sufficient hardenability, appropriate martensite start temperature (Ms), and austenite finish temperature (Ac3) through alloy design, enabling components to attain optimal mechanical properties.

[0080] In this invention, the low-carbon design endows martensite with excellent toughness, while maximizing martensitic structure achieves high strength. Specifically, a fully martensitic microstructure ensures higher yield strength in components, effectively preventing premature deformation failure during collisions. To this end, the invention optimizes the content of cost-effective alloy elements (Mn, Si, and Cr) to: Reduce non-martensitic phases (e.g., ferrite, bainite). Ensure sufficient hardenability via the formula: Mn+0.26Si+1.3Cr≥2.20%. Preferred Range: Mn+0.26Si+1.3Cr≥2.25%. Further Optimization: Mn+0.26Si+1.3Cr≥2.50% for enhanced hardenability, enabling near-full martensitization.

[0081] The microstructure of the component according to the present invention consists of the following area percentages: Bainite: <5%, Austenite (retained): <3%, Ferrite: <3%, The remainder being martensite. Preferably, the combined content of bainite, ferrite, and retained austenite is <5%. More preferably, their total content does not exceed 3%.

[0082] Furthermore, to ensure the safety of automotive components, it is essential to achieve both high strength and high toughness. It is well known that during steel heat treatment, phase transformation latent heat is generated as martensite forms. The newly formed martensite undergoes self-tempering as this latent heat is released, which alleviates transformation stresses and reduces the interaction between carbon (C) and dislocations in martensite, ultimately improving the toughness of the martensitic matrix. During component hot deformation and mold cooling, the martensite start (Ms) temperature plays a critical role in controlling martensite formation and self-tempering effects. Those skilled in the art understand that carbon (C) significantly lowers the Ms temperature in steel. Therefore, in medium-to-high carbon martensitic steels, alloy design is typically employed to raise the Ms temperature, thereby enhancing self-tempering effects and improving the toughness of inherently hard and brittle high-C martensite. However, in the present invention's low-C steel, the interaction between carbon and dislocations in martensite is far weaker than in high-C steels. Consequently, the intrinsic toughness of martensite is better, and the self-tempering effect has limited impact on toughness improvement. Additionally, the low-C content already results in a relatively high Ms temperature, leading to excessive self-tempering, martensite temper decomposition softening, and cementite precipitation, which significantly degrade both strength and toughness. Therefore, in the context of this invention's low-C design, high Mn content combined with optimized Si and Cr ratios is employed to control the Ms temperature within 350-410°C, preferably 360-400°C, thereby moderating the self-tempering extent and achieving further optimization of strength-toughness balance.

[0083] Furthermore, reducing the carbon content significantly increases the austenite finish temperature (Ac3), making it impossible to achieve a fully austenitic structure under conventional hot stamping heating conditions. To ensure complete austenitization, this invention employs high Mn content combined with optimized ratios of Si and Cr to control the Ac3 temperature ≤900°C.

[0084] Preferably, this invention aims to control the precipitation of microalloy carbides (e.g., Nb, V, Ti) in the microstructure. An appropriate amount of microalloy carbides enhances grain refinement and precipitation strengthening, thereby improving the strength-toughness balance of components. However, excessive carbide precipitation consumes carbon in martensite, reducing the interaction between carbon and dislocations, which is detrimental to achieving high strength. Therefore, the contents of Nb, V, and Ti must be strictly limited within the aforementioned ranges.

[0085] The hot-stamped component of this invention may additionally be coated with a metallic coating layer, which can be based on aluminum-based alloy coatings or zinc-based alloy coatings.

[0086] It should be noted that the above design is also applicable to the steel plates used for manufacturing the aforementioned hot-stamped forming components.

[0087] The high strength of the material enhances the lightweight performance of hot-stamped forming components. During collision failure, components first undergo local plastic deformation at the impact site, followed by bending and cracking in the deformed region. Therefore, the yield strength becomes a critical evaluation criterion, as it represents the critical stress at which steel transitions from elastic to plastic deformation. A high yield strength significantly delays the onset of plastic deformation, preventing premature failure and facilitating post-collision repairs. The bending performance of the material defines its foldability, where superior bendability delays fracture initiation during bending deformation and avoids early brittle failure. This allows the component to sustain deformation and absorb collision energy, significantly improving energy absorption efficiency. Additionally, a higher peak load at the point of bending fracture indicates better energy absorption performance. Given the variability of vehicle collision scenarios, evaluating crash safety solely based on tensile strength (from tensile tests) and maximum bending angle (from VDA bending tests, influenced by material thickness) is insufficient. To comprehensively assess crash safety, the inventors propose incorporating peak force and bending energy absorption into the evaluation criteria.

[0088] The commonly used methods to evaluate material toughness is the static three-point bending test (VDA bending test, VDA 238-100 standard). Through testing, parameters such as maximum bending angle, peak force at maximum bending angle, bending energy absorption (W), and fracture strain can be obtained, which collectively reflect the material's ability to resist bending deformation failure. The room-temperature tensile test (GB / T 228.1 standard) is a common method for measuring material strength. The yield strength and tensile strength reflect the material's ability to resist deformation and failure under tensile loading. Furthermore, during the VDA bending test (VDA 238-100), the bending moment (M) in the deformation zone is proportional to the square of the plate thickness (t 2< ). The bending energy absorption (W) represents the total energy required to induce bending deformation throughout the test, which is also proportional to t 2< . To eliminate the influence of thickness differences on energy absorption evaluation, the parameter W / t 2< is used to characterize the unit bending energy absorption performance.

[0089] Based on the above, the components of this invention exhibit the following mechanical properties: Yield strength (YS): 880-1100 MPa Tensile strength (TS): ≥1000 MPa Elongation: ≥5% Fracture strain (ε): ≥0.6 Preferred ranges: Yield strength (YS): 930-1050 MPa Tensile strength (TS): ≥1060 MPa Elongation: ≥7% Fracture strain (ε): ≥0.7 Additionally, the bending energy absorption in static three-point bending tests meets: W≥50, 000mm* N, W / t2≥2.88×104 N / mm(normalized energy absorption per unit thickness).

[0090] The present invention provides a hot-stamping method for producing the aforementioned hot-stamped components, comprising: A) Supply a blank or preform obtained by preforming the hot-stamping steel plate. B) Full Austenitization Step:Heat the blank or preform to≥900°C and hold at this temperature. Preferred Conditions: The heating is performed in a furnace with a maximum temperature range of 900-980°C to achieve full austenitization. To ensure uniform heating across different thicknesses, the total heating time (T, in minutes) in the furnace satisfies: T=t~ (t×3+10) minutes where t is the thickness of the blank / preform (in mm).C) Hot Stamping and Forming Step

[0091] Transfer the heated blank / preform to a press for hot deformation, followed by cooling in a mold to <300°C to obtain the hot-stamped component. Cooling Requirements: In non-deformed regions, the average cooling rate from 700°C to Ms temperature is≥40°C / s. In deformed regions, cooling is accelerated to achieve higher average cooling rates. Preferred Ms Temperature Range: 350°C≤Ms temperature≤410°C.

[0092] A high average cooling rate minimizes the formation of non-martensitic phases (e.g., bainite, ferrite, retained austenite), thereby ensuring the strength and toughness of the hot-stamped component.

[0093] To further optimize the process, since hot deformation promotes ferrite formation, higher cooling rates are required to suppress ferrite transformation as deformation increases. Thus, in Step C, the following constraints apply:C) Hot Stamping and Forming Step (Optimized)

[0094] In at least one hot-deformed region of the component: Deformation amount≤10%.Average cooling rate from 700°C to Ms temperature≥50°C / s.

[0095] More preferably: Deformation amount≤20% in at least one region. Average cooling rate from 700°C to Ms temperature≥60°C / s.

[0096] The following will describe the present invention in further detail with reference to exemplary embodiments. The embodiments or experimental data provided below are intended to illustratively demonstrate the invention, and it should be understood by those skilled in the art that the invention is not limited to these embodiments or experimental data.

[0097] The steel plate with the composition shown in Table 1 is prepared through the following manufacturing process: a) Steelmaking: The steel is smelted using a vacuum induction furnace, electric furnace, or converter according to the composition listed in Table 1. b) Hot Rolling: The slab is heated to 1200°C and held for 2 hours. Hot rolling is performed at 800-1200°C, followed by coiling below 700°C to form a hot-rolled coil. The hot-rolled coil is pickled to remove oxide scale generated during rolling. c) Cold Rolling: The pickled hot-rolled coil undergoes cold rolling with a reduction rate of 30%-70% to achieve a final thickness of 1.4-1.9 mm (cold-rolled coil). d) Hot Coating: For T1, T2, T4, T5, T7, T8, T9, T10, T11, and CT1-CT3, apply hot-dip aluminum alloy coating to obtain the final test steel plate. Single-side coating weight:20-75 g / m 2< .For T6, apply hot-dip zinc alloy coating to obtain the final test steel plate. Single-side coating weight: 40-80 g / m 2< . e) Annealing: For T3, perform annealing on the cold-rolled coil to obtain the final test steel plate. Note: T3 does not undergo hot coating. Table 1. Chemical composition of the base steel (wt.%, the remainder being Fe and inevitable impurities). CSiMnCrAlTiNbBVPSOther ElementsMsAc 3Mn + 0.26S i+ 1.3CrSi + CrT 10.0680.171.900.200.03 80.0280.00250.0150.0 083968832.200.3 7T 20.070.302.020.080.02 70.0310.00290.0140.0 093938852.200.3 8T 30.0720.111.930.270.02 70.0310.050.00290.0130.0 093928772.310.3 8T 40.0590.192.030.150.040.0300.03, 0.030.00350.0140.0 073948902.270.3 4T 50.0630.212.180.230.040.0460.00350.0 50.0110.0 053838842.530.4 4T 60.0670.252.230.220.050.0180.00150.1 10.0120.0 093798942.580.4 7T 70.0890.202.400.110.01 50.0330.00250.0130.0 083678672.600.3 1T 80.070.112.470.250.02 50.020.00350.0120.0 073668672.820.3 6T 90.0690.101.900.310.03 70.0340.00270.0130.0 09Mo: 0.1, Sb: 0.006, Ca: 0.0023928802.310.4 1T 1 00.0610.262.080.150.03 50.0250.030.00320.0140.0 06Ni: 0.1, Sn: 0.01, REM: 0.0023898842.340.4 1T 1 10.0870.222.280.180.01 30.0310.00270.0150.0 08Cu: 0.1, Co: 0.1, Mg: 0.0023728702.570.4 0C T 10.070.421.540.040.040.0220.03 80.00210.0130.0 094179001.700.4 6C T 20.110.202.720.130.050.030.00150.0160.0 073448642.940.3 3C T 30.0690.301.820.220.110.0310.00290.0160.0 063999062.180.5 2

[0098] T1-T11 are exemplary steels within the compositional scope of the present invention that satisfy the Mn, Si, and Cr ratio relationships. CT1-CT3 are comparative steels outside the scope of the present invention.

[0099] T1-T11 have the following compositional and thermal properties: Carbon (C) content: 0.059-0.089 wt.%. Mn + 0.26Si + 1.3Cr≥2.20%. Si + Cr≤0.47%. Calculated Ac3 temperature: ≤900°C (all T1-T11). Calculated Ms temperature: 366-396°C (all T1-T11).

[0100] As a comparison, CT1 has a higher Si content and lower Mn content, resulting in a calculated Mn + 0.26Si + 1.3Cr value below 2.20, which indicates poor hardenability. Additionally, the high Si content promotes ferrite formation, facilitating the generation of non -martensitic phases, thereby hindering strength improvement. Furthermore, the calculated Ms temperature is relatively high (417°C), preventing effective control of martensite self-tempering and weakening strength enhancement. For CT2, the elevated C and Mn contents lead to a lower Ms temperature, insufficient for effective martensite self-tempering to improve toughness. Although CT3's composition falls within the scope of the present invention, its calculated Mn + 0.26Si + 1.3Cr value remains below 2.20, indicating insufficient hardenability. In this case, even though Si is within the specified range, the hardenability is inadequate to suppress Si's promotion of ferrite formation, resulting in minor ferrite segregation that compromises strength. Moreover, the calculated Ac3 temperature is relatively high (906°C), making it difficult to achieve complete austenitization during hot stamping heating. Finally, CT3 contains elevated levels of both Si and Cr, with a combined Si + Cr value of 0.52%. During the preparation of CT3's coated steel plate samples, more severe plating defects (e.g., incomplete coating coverage and poor adhesion) were observed compared to other examples. To further improve the surface quality of coated steel plates and ensure more uniform microstructures and enhanced performance of the steel plate after hot stamping, the Si + Cr value should not exceed 0.50%.Critical Cooling Rate Test

[0101] The critical cooling rate (CCD) for complete martensitic transformation of T1, T2, T4, T5, and CT1 compositions was tested using the DIL805 dilatometer. Generally, a higher CCD indicates a greater cooling requirement for full martensitic transformation, implying poorer hardenability of the material. Initially, specimens were heated to 930°C and held at this temperature. After cooling at rates ranging from 20-70°C / s (with and without applying 10% thermal deformation) to below 200°C, the expansion curves during the cooling stage (700-400°C) were analyzed. The presence of non-martensitic transformations in this temperature range was evaluated by identifying inflection temperatures. The minimum cooling rate without non-martensitic transformation was recorded as the critical cooling rate for each material, as shown in Table 2. In conventional hot stamping processes, thermal deformation typically does not exceed 10%. Only components requiring deep drawing may involve deformation up to 20%. Therefore, 10% deformation was adopted as the baseline for determining the critical cooling rate for martensitic transformation in this study. Additionally, as presented in Table 1, the calculated values of Mn + 0.26Si + 1.3Cr for T1 and T2 are both 2.20. Compared to other steels, such as T4 (2.27) and T5 (2.53), T1 and T2 exhibit lower hardenability. Since T1 and T2 represent the materials with the most stringent cooling requirements, their CCD results are used as the benchmark. If these compositions meet the cooling criteria, other steels with higher Mn + 0.26Si + 1.3Cr values will inherently satisfy the requirements. Table2:Test Results of Critical Cooling Rates (CCD) for Experimental SteelsNumb erHot - deformatio n amountCritical cooling rate, °C / sNon - martensitic phase transformation start temperature, °C (marked as "×" if there is no non - martensitic phase transformation)20°C / s25°C / s30°C / s40°C / s50°C / s70°C / sT50<20××××××T420~25550×××××T125~30462××××T225~30473××××CT140~50552××T510%25~30492××××T440~50470××T140~50486××T250~60467×CT1>70540

[0102] As shown in Table 2, under non-deformation conditions, the critical cooling rate (CCD) for T5 is below 20°C / s, T4 is approximately 25°C / s, while T1 and T2 are around 30°C / s. Notably, the critical cooling rate for 22MnB5, a material widely used in hot stamping production, is approximately 27°C / s. This indicates that through optimized design of Mn, Si, and Cr contents (with Mn + 0.26Si + 1.3Cr ≥ 2.20%), the steel in this invention achieves sufficient hardenability even with low carbon content. Consequently, under existing hot stamping production conditions, full martensitic structures can be reliably obtained during cooling.

[0103] In contrast, due to the Mn+0.26Si+1.3Cr value of approximately 1.70%, CT1 still exhibits significant non-martensitic phase transformations during cooling even at a cooling rate of 40°C / s under non-deformation conditions, indicating poorer hardenability. Consequently, under existing hot stamping conditions, CT1 is likely to form excessive ferrite and bainite, which would adversely affect production control and compromise the strength assurance of the final part.

[0104] Under 10% thermal deformation conditions, the critical cooling rates (CCD) for T5 and T4 fall within 25-30°C / s and 40-50°C / s, respectively, while T2 exhibits a CCD of 50-60°C / s. This indicates that the lower the Mn+0.26Si+1.3Cr value, the higher the required CCD. Notably, T1 and T2 share the same Mn 0.26Si+1.3Cr value, yet T1 demonstrates a lower CCD (40-50°C / s) compared to T2. This discrepancy arises because T1 contains higher Cr content (favorable for suppressing bainite formation) and lower Si content (less promoting of ferrite), which collectively reduce the CCD. Therefore, even with comparable hardenability under non-deformation conditions, the combination of higher Cr and lower Si enhances the ability to achieve full martensitic structures under thermal deformation conditions.

[0105] In contrast, even at a cooling rate of 70°C / s, CT1 steel still exhibits significant non-martensitic phase transformations. This is primarily attributed to its insufficient hardenability, excessively high Si content, and insufficient Cr content.

[0106] In light of the above, achieving a uniform martensitic structure requires a higher critical cooling rate (CCD) under thermal deformation conditions compared to non-deformation scenarios. With 10% deformation as the baseline, for the steel in this invention, ensuring an average cooling rate above 60°C / s allows the component to attain a uniform martensitic structure. Preferably, by increasing Cr content (above 0.11%) and reducing Si content (below 0.30%), the required average cooling rate can be further lowered to above 50°C / s while still achieving a uniform martensitic structure. In contrast, for the CT1 sample, even with an average cooling rate exceeding 70°C / s, significant ferrite and / or bainite formation persists, preventing the component from achieving optimal performance. Therefore, the steel in this invention optimizes Mn, Si, and Cr ratios to enhance hardenability, suppresses non-martensitic phase formation, and reduces production control requirements while maintaining performance.Testing the properties of hot stamping-formed components.

[0107] Hot stamping of hat-shaped components was conducted using steel plate specimens with varying thicknesses from T1 to T11 and CT1 to CT3. Specimens measuring 230×550 mm were heated to 910°C and held at this temperature to achieve complete austenitization. Subsequently, the specimens were transferred to a press for hot stamping within a transfer time of 10-12 seconds. Hot stamping was then performed in the press, followed by holding and cooling for 8-10 seconds until the temperature dropped below 250°C. The hat-shaped components were then removed, with the average cooling rate during this process reaching over 50°C / s. Samples were taken from areas of the component sidewalls that experienced thermal deformation. The microstructures of the post-hot-stamping specimens were examined under an optical microscope and scanning electron microscope (SEM). Mechanical properties, including tensile strength, elongation, maximum bending angle, and peak force, were tested in accordance with GB / T 228.1 (room temperature tensile testing standard) and VDA-238 (three-point bending standard). Additionally, the bending energy absorption of the specimens prior to bending failure was calculated based on the load-displacement curve from the VDA-238 three-point bending test. This was determined by integrating the area under the curve up to the peak load. To minimize measurement errors, all reported results represent the average of three independent tests.

[0108] The following steps outline the procedure for determining the bending fracture strain of component specimens:(1)Conduct a static three-point bending test to establish the peak bending angle (αpeak) of the specimen.(2)Based on the experimental results, select at least three interrupted bending angles (αL)-the bending angle of the specimen under load-ensuring that αL ≥ 50% αpeak.(3)Stop loading when the specimen reaches the selected αL, then measure the unloaded bending angle (αUL) after unloading.(4)Place the unloaded specimen under an optical microscope and measure the inner (Ri) and outer (Ro) surface radii of the most severely deformed region.(5)Use Equation (1) to compute the equivalent (plastic) strain (ε) at the outer surface of the most severely deformed region for each αUL. This represents the equivalent strain of the specimen at the selected αL during bending, thereby establishing the ε-αL relationship.(6)Based on the fitted ε-αL curve, extrapolate to determine the bending fracture strain (εfracture)-i.e., the equivalent strain at αL = αpeak. ε = 2 3 ⋅ ln R o R i The final results of the microstructure, tensile properties, and VDA bending performance are presented in Table 3. Table 3 test results of microstructure and properties for each specimen.Numb erPlate Thic knes s t / mmSurface ConditionComposition of Non - Martensitic Phases (by Area Percentage)YS MPaTS MPaElongatio n %Maxim um Bendin g Angle / °εPeak Force NW mm*NW / t 2< ×10 4< T11.41Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%92610717.897.50.827674629063.16T21.39Aluminu m-PlatedBainite < 5%, Austenite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%88010547.695.60.817532603723.12T31.41UncoatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%93811057.589.40.737902576682.9T41.56Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%93210707.295.10.829261715542.94T5155Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%95411207999.70.879725778193.24T61.43Zinc - PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%97911167.196.20.818239655453.21T71.78Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%99511628890.7913120963453.04T81.41Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%104111867.588.30.78241605203.04T91.39Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%93410827.696.70.817850632743.27T101.54Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%94510657.196.10.819175731813.09T111.81Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%100211787.289.50.7712973952262.91CT11.42Aluminu m - PlatedBainite ≥ 10%, Austenite + Ferrite < 5%, Nb - V - Ti Micro - alloyed Carbides < 0.2%8169617.51040.96568572402.84CT21.41Aluminu m-PlatedAustenite + Bainite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%96712676.279.30.598545570792.87CT31.42Aluminu m - Plated5% < Bainite < 10%, Austenite + Ferrite < 3%, Nb - V - Ti Micro - alloyed Carbides < 0.2%85810057.398.60.836905567922.82

[0109] Figure 1 shows the typical microstructure of the T1 component specimen of the present invention. In its microstructure, the martensite phase is dominant, and the total amount of non - martensitic phases is less than 3% (by area percentage). The typical microstructure of the T1 component specimen is applicable to the T3 - T11 component specimens. Figure 2 shows the typical microstructure of the T2 component specimen of the present invention, in which there is a small amount of bainite, accounting for less than 5% by area percentage. Figure 3 shows the typical microstructure of the CT1 component specimen, in which there is a relatively large amount of bainite, accounting for more than 10% by area percentage. Combined with the microstructures of various component specimens and the values of Mn+0.26Si+1.3Cr, it can be confirmed that when Mn+0.26Si+1.3Cr≥2.20%, a non - martensitic phase composition of less than 8% by area percentage can be obtained, indicating improved hardenability, which is beneficial to improving the strength of the steel. In contrast, the microstructures of CT1 and CT3 show that when Mn+0.26Si+1.3Cr<2.20%, there is a non - martensitic phase composition of more than 8% by area percentage, indicating that the hardenability of the steel is insufficient and cannot effectively improve the strength of the steel.

[0110] As shown in Table 3, the yield strength of the component specimens T1 - T11 of the present invention ranges from 880 to 1041 MPa, the tensile strength ranges from 1054 to 1186 MPa, the elongation is in the range of 7% - 8%, the maximum bending angle reaches 88°-99.7°, the fracture strain is above 0.7, and the W / t 2< value is not less than 2.9×10 4< N / mm, demonstrating an excellent match between strength and toughness.

[0111] Compared with the T1 component specimen, although the maximum bending angle and fracture strain of the CT1 component specimen have improved, there are significant reductions in yield strength, tensile strength, peak force, and W / t 2< value, which are 816 MPa, 961 MPa, 6568 N, and 2.84×10 4< N / mm respectively, failing to meet the usage requirements. This is because the microstructure of the CT1 component specimen contains more than 10% bainite and a small amount of ferrite. As previously mentioned, bainite can improve the toughness of the material. Therefore, the CT1 component specimen achieves a maximum bending angle of 104° and a fracture strain of 0.90. However, the presence of a relatively large amount of bainite and a small amount of ferrite is detrimental to the strength of the material. In addition, the high Ms temperature causes excessive self - tempering of martensite, which also leads to a decrease in strength. The CT1 component specimen has a microstructure with a relatively large amount of bainite and a small amount of ferrite for the following reasons: On the one hand, the composition of the CT1 component specimen is outside the scope of the present invention. Specifically, the low Mn content and the inappropriate ratio of Mn, Si, and Cr (Mn+0.26Si+1.3Cr=1.70) result in poor hardenability, tending to form non - martensitic structures such as ferrite and bainite. On the other hand, under the condition of poor hardenability, the low Cr content (0.04%) is insufficient to suppress the formation of bainite, and the high Si content (0.42%) promotes the formation of ferrite, further leading to the formation of a relatively large amount of bainite and ferrite in the microstructure under the same hot stamping forming conditions.

[0112] The CT2 component specimen exhibits high strength, with a tensile strength of 1267 MPa, and features a microstructure similar to that of the T1 component specimen. This is primarily due to the fact that Mn + 0.26Si + 1.3Cr = 2.94, indicating that the CT2 component specimen possesses high hardenability, ensuring a fully martensitic structure. Additionally, the relatively high carbon (C) content contributes to this strength. However, the toughness of the CT2 component specimen is lower than that of all other component specimens, with a maximum bending angle of only 79.3°and a fracture strain of 0.59. This is because a manganese (Mn) content exceeding 2.70% can lead to microsegregation, thereby reducing the ductility and toughness of the steel. Moreover, the slightly elevated C content in the CT2 component specimen, combined with a lower Ms temperature, results in insufficient martensite self-tempering during cooling. This insufficient self-tempering fails to effectively reduce the interaction between carbon and dislocations, ultimately preventing significant improvement in martensite toughness. These two factors collectively result in the insufficient toughness of the CT2 component specimen. Consequently, even if the hardenability requirements are met, if the composition exceeds the scope of the present invention, it is not possible to achieve components that simultaneously possess high strength and high toughness.

[0113] Compared with the T1 component specimen, the CT3 component specimen exhibits good toughness. However, its yield strength decreases by approximately 7% and fails to reach 880 MPa. This is because, although the compositions of both specimens fall within the scope of the present invention, the ratio of Mn, Si, and Cr in the CT3 component specimen does not meet the requirement of Mn+0.26Si+1.3Cr≥2.20%. As a result, its hardenability is relatively poor, leading to the formation of a relatively larger amount of bainite (with an area percentage of approximately 5 - 10%). Consequently, its yield strength is lower.

[0114] Based on the above analysis, to achieve the desired combination of high strength and toughness-specifically, a tensile strength of no less than 880 MPa, a maximum bending angle of no less than 80°, a fracture strain of no less than 0.6, and a W / t 2< value of no less than 2.88×10 4< N / mm-the chemical composition must satisfy two critical conditions: 1.All alloying elements must fall within the compositional range defined by the present invention. 2.The combined content of Mn, Si, and Cr must meet or exceed the threshold of Mn+0.26Si +1.3Cr≥2.20%. Neither condition alone is sufficient; both must be simultaneously fulfilled.

[0115] Furthermore, as shown in Table 3, overall, when the toughness requirements are met-namely, a maximum bending angle of no less than 80°, a fracture strain (ε) of no less than 0.6, a peak force of no less than 7500 N, and a bending energy absorption of W / t 2< ≥2.88×10 4< N / mm-the strength of the T1-T11 component specimens exhibits an overall increasing trend with the rise in the value of Mn+0.26Si+1.3Cr, indicating enhanced hardenability. In particular, the T4 component specimen, with an Mn+0.26Si +1.3Cr value of 2.27%, achieves a yield strength of 932MPa. Therefore, when the Mn+0.26Si+1.3Cr value is 2.25% or greater, a yield strength of no less than 930MPa can be obtained. Furthermore, the T5 component specimen, with an Mn+0.26Si+1.3Cr value of 2.53%, achieves a yield strength of 954MPa. Thus, when the Mn+0.26Si+1.3Cr value is 2.50% or greater, a yield strength of no less than 950MPa can be achieved.

[0116] The toughness of the T1 component specimen is similar to that of the T2 component specimen; however, the yield strength and tensile strength of the T2 component specimen are significantly lower than those of the T1 component specimen. This is because, although both specimens have comparable carbon (C) contents and identical values of Mn+0.26Si+1.3Cr, the Cr content in the T2 component specimen is lower than that in the T1 component specimen. Consequently, the Cr content in the T2 specimen is insufficient to suppress bainite formation during cooling, resulting in a microstructure containing a small amount of bainite (as shown in Figure 2). The presence of bainite reduces the yield strength and tensile strength of the T2 component specimen. Similarly, the CT1 component specimen also exhibits a substantial amount of bainite due to its low Cr content, leading to a significant decrease in both yield strength and tensile strength. Therefore, while ensuring hardenability and toughness, increasing the Cr content can further enhance the strength. Preferably, the Cr content should be maintained within the range of 0.11%≤Cr≤0.5%, and more optimally, within 0.12%≤Cr≤0.4%.

[0117] Additionally, the T3 and T6 component specimens, representing uncoated and galvanized samples respectively, both meet all performance requirements. This demonstrates that hot-stamped components, whether galvanized, aluminized, or uncoated, can satisfy the expected performance criteria. Furthermore, the experimental results of T1-T11 specimens confirm that components with varying thicknesses designed according to this invention can consistently meet the anticipated performance standards.

[0118] The aforementioned results demonstrate that, within the matrix steel composition range of the present invention and satisfying the condition of Mn+0.26Si+1.3Cr≥2.20%, hot-stamped components with either galvanized, aluminized, or uncoated coatings can all meet the expected performance requirements, exhibiting a combination of high strength, high toughness, and superior energy absorption capacity, thereby enhancing usage safety. Furthermore, performance can be further improved by more precisely controlling the range of alloying elements.

[0119] The above embodiments and experimental data are provided solely for exemplary illustration of the present invention. It should be apparent to those skilled in the art that the present invention is not limited to these specific examples and that various modifications can be made without departing from the scope of protection defined by the present invention.

Claims

1. A hot stamping steel sheet, wherein the base steel of the hot stamping steel sheet comprises, by mass percentage: 0.053%≤C≤0.10%, 0.05%≤Si≤0.30%, 1.81%≤Mn≤2.7%, 0.01%≤Cr≤0.7%, 0.01%≤Al≤0.5%, 0.0005%≤B≤0.005%, 0.015%≤Ti≤0.05%, 0≤Nb+V≤0.2%, 0.001%≤P≤0.100%, 0.0001 % ≤S≤0.100%, Fe≥95% and inevitable impurities, the contents of Mn, Si, and Cr satisfy the condition: Mn+0.26Si+1.3Cr≥2.20%.

2. According to the hot stamping steel sheet described in claim 1, the steel matrix further comprises, by mass percentage, at least one of the following elements: 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01% and 0.0001%≤REM≤0.01%.

3. According to the hot stamping steel sheet described in claim 2, wherein the total content of the following elements satisfies: 0.0001%≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30%.

4. According to the hot stamping steel sheet described in claim 1, the remainder of the steel matrix consists of iron (Fe) and inevitable impurities, by mass percentage.

5. According to the hot stamping steel sheet described in claim 3, the remainder of the steel matrix consists of iron (Fe) and inevitable impurities, by mass percentage.

6. According to the hot stamping steel sheet described in claim 1, wherein the sum of Mn, Si, and Cr contents satisfies: Mn + 0.26Si + 1.3Cr ≥ 2.25%.

7. According to the hot stamping steel sheet described in claim 1, wherein Mn + 0.26Si + 1.3Cr ≥ 2.50%.

8. A hot stamping steel sheet according to any one of claims 1 to 7, wherein the chromium (Cr) content satisfies: 0.11% ≤ Cr ≤0.5%.

9. A hot stamping steel sheet according to any one of claims 1 to 7, wherein 0.12% ≤ Cr ≤ 0.4%.

10. A hot stamping steel sheet according to any one of claims 1 to 7, wherein:2.01% ≤ Mn ≤ 2.50%, and / or 0.055% ≤ C ≤ 0.09% and 0.10% ≤ Si ≤ 0.26%.

11. A hot stamping steel sheet according to any one of claims 1 to 7, wherein the total content of niobium (Nb) and vanadium (V) satisfies: 0.01% ≤ Nb + V ≤ 0.2%.

12. A hot stamping steel sheet according to any one of claims 1 to 7, wherein the martensitic transformation start temperature (Ms) of the steel sheet is calculated based on the alloying elements in mass percentage using the formula: Ms=520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu and satisfies:350°C≤Ms≤410°C.

13. A hot stamping steel sheet according to any one of claims 1 to 7, wherein the austenite transformation finish temperature (Ac3) of the steel sheet is calculated based on the mass percentages of alloying elements using the following formula: Ac3=912-250C-16Mn+48Si-2Cr-16Ni+95V+96Ti+210Al-10Cu and satisfies: Ac3 ≤ 900°C.

14. A hot stamping steel sheet according to any one of claims 1 to 7, wherein the sum of silicon (Si) and chromium (Cr) contents satisfies: Si + Cr ≤ 0.47%.

15. A hot stamping formed component, the steel sheet matrix of the hot stamping formed component comprises, by mass percentage: Carbon (C): 0.053% ≤ C ≤ 0.10%, Silicon (Si): 0.05% ≤ Si ≤ 0.30%, Manganese (Mn): 1.81% ≤ Mn ≤ 2.7%, Chromium (Cr): 0.01% ≤ Cr ≤ 0.7%, Aluminum (Al): 0.01% ≤ Al ≤ 0.5%, Boron (B): 0.0005% ≤ B ≤ 0.005%, Titanium (Ti): 0.015% ≤ Ti ≤ 0.05%, Niobium + Vanadium (Nb + V): 0 ≤ Nb + V ≤ 0.2%, Phosphorus (P): 0.001% ≤ P ≤ 0.100%, Sulfur (S): 0.0001% ≤ S ≤ 0.100%, Iron (Fe): ≥ 95% and inevitable impurities, wherein the contents of Mn, Si, and Cr satisfy: Mn + 0.26Si + 1.3Cr ≥ 2.20%.The microstructure of the steel sheet matrix, by area percentage, includes: bainite: < 5%, austenite: < 3%, ferrite: < 3%, Nb-V-Ti microalloyed carbides: < 0.2%, The remainder consists of martensite.

16. A hot stamping formed component according to claim 15, the steel sheet matrix of the hot stamping formed component further comprises, by mass percentage, at least one of the following elements: 0.01%≤W≤0.30%, 0.01%≤Mo≤0.30%, 0.01%≤Ni≤0.30%, 0.01%≤Cu≤0.30%, 0.01%≤Co≤0.30%, 0.005%≤Sn≤0.30%, 0.005%≤Sb≤0.100%, 0.0001%≤Ca≤0.01%, 0.0001%≤Mg≤0.01%, 0.0001%≤Zr≤0.01% and 0.0001%≤REM≤0.01%.wherein the total content of the above elements satisfies: 0.0001%≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤ 0.30%.

17. A hot stamping formed component according to claim 15, wherein the remainder of the steel sheet matrix consists of iron (Fe) and inevitable impurities, by mass percentage.

18. A hot stamping formed component according to any one of claims 15 to 17, wherein the sum of bainite, ferrite, and retained austenite in the microstructure of the steel sheet matrix does not exceed 3% by area percentage.

19. A hot stamping formed component according to any one of claims 15 to 17, wherein the component has: a yield strength (YS) of not less than 930-1050 MPa, a tensile strength (TS) of not less than 1060 MPa, an elongation of not less than 7%, a fracture strain (ε) of not less than 0.7, and a bending energy absorption per unit thickness (W / t2) during static three-point bending that satisfies: W / t2 ≥ 2.88×104 N / mm.

20. A hot stamping formed component according to any one of claims 15 to 17, wherein the component is made from a hot stamping steel sheet according to any one of claims 2 to 14.

21. A hot stamping method for producing a hot stamped component according to any one of claims 15 to 20, comprising: A) Providing a blank or preform made from a hot stamping steel sheet according to any one of claims 1 to 14, wherein: the blank is an unformed material, or the preform is obtained by preliminary shaping of the hot stamping steel sheet; B) Full austenitization treatment: heating the blank or preform from step A to above 900°C and holding at this temperature; C) Hot stamping forming treatment: after completing step B, transferring the heated blank or preform to a press for hot deformation, followed by cooling in the die to below 300°C to obtain the hot stamped component, wherein: the average cooling rate from 700°C to Ms temperature in non-deformed regions is not less than 40°C / s, and the average cooling rate in deformed regions is higher than that in non-deformed regions.

22. The method according to claim 21, wherein in step C: at least one region of the blank or preform undergoes a hot deformation amount of not more than 10%, and the average cooling rate from 700°C to Ms temperature in said at least one region is not less than 50°C / s.

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