Steel sheet manufacturing method, hot stamping formed component, and hot stamping steel sheet

Abstract

The present invention relates to a steel sheet for hot stamping, wherein a steel sheet substrate of the steel sheet for hot stamping contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95%, and inevitable impurities, wherein Mf of the steel sheet substrate satisfies Mf≧230°C; and hardenability factor h satisfies h≧8.0.

Description

[Technical Field]

[0001] The present invention relates to a steel sheet for hot stamping, a hot stamped component and a method for manufacturing said steel sheet. [Background technology]

[0002] Weight reduction is an important approach to achieving energy conservation and exhaust emission reduction in the automotive industry. The strength of hot stamped components, obtained by hot deformation and cooling of hot stamping steel, can exceed 1500 MPa. When applied to the car body, the hot stamped components can achieve significant weight reduction and ensure the collision safety of the vehicle. Therefore, the adoption of hot stamped components in automobiles has been steadily increasing.

[0003] However, improving the strength of steel generally reduces its plasticity and toughness. Therefore, in automotive applications, engineers are increasingly concerned about issues resulting from insufficient toughness and delayed cracking in hot-stamped components. For example, during actual production processes, some hot-stamped components often develop brittle cracks during placement, processing, transportation, or welding after hot stamping, leading to scrapping of parts or entire vehicle bodies and increased production costs. Even more seriously, some hot-stamped components develop delayed cracks after vehicle assembly, posing a safety risk to the vehicle while it is in motion. Therefore, ensuring sufficient toughness and delayed cracking resistance is crucial, especially when using hot-stamped components with strengths exceeding 1700 MPa.

[0004] Chinese Patent No. 106399837 provides a steel or formed component with high toughness and delayed crack resistance. Adding more than 0.11% V precipitates a large amount of nanoscale VC and / or (V, Ti, Nb)C complex carbides in the microstructure of the formed component. The precipitation of these carbides refines the grains and improves the yield strength and tensile strength of the material. Furthermore, consuming carbon in austenite reduces the proportion of brittle twinned martensite, improving the toughness of the martensite and thereby improving the material's resistance to hydrogen-induced delayed cracking. To achieve a fully martensite structure, this patent also adds alloying elements such as Mn and Mo to improve hardenability. However, such high-alloy designs make strip steel prone to the formation of hard, brittle martensite during processes such as hot rolling, annealing, and coating, which increases the likelihood of cracking and strip breakage during cold rolling. Furthermore, the relatively high strength of the annealed or coated material makes it difficult to process and control the plate shape throughout the manufacturing process, significantly increasing the difficulty and cost of steel plate production.Furthermore, this patent does not take into account the effect of coarse TiN particles on the toughness of the steel plate.

[0005] In Chinese Patent No. 108374127, to overcome the effect of TiN particles on the toughness of steel plates, the following two Ti-free alloy compositions are proposed: 1. Low B content design (the effect of B is replaced by a reasonable proportion of hardenable alloy elements such as Mn, Cr, and Mo). That is, when B≦0.0005%, 29 * Mo+16 * Mn+14 * Cr+5.3 *Ni ≥ 30% is satisfied. 2. High B content design (Adding a certain amount of Al to replace the bond between Ti and N, avoiding the precipitation of large-sized TiN particles, and improving the toughness of the hot stamping steel sheet and its hot stamp forming components. That is, when 0.0005% < B ≤ 0.005%, 0.4 - 1.0% of Al is included.) was designed. The hot stamp forming component obtained by the above alloy design has a yield strength of 1200 - 1800 MPa, a tensile strength of 1500 - 2150 MPa, an elongation after annealing of 7 - 10%, and an impact toughness at -40°C ≥ 45 J·cm -2 A 99.5% confidence level can be achieved. However, in the specification of Chinese Patent No. 108374127, the addition of high hardenability elements such as Mn, Cr, and Mo is not considered, so it is easy to generate hard and brittle martensite in the manufacturing process, which is disadvantageous to processes such as cold rolling, annealing, coating, and subsequent steel processing. Furthermore, when 0.4 - 1.0% of Al is added, the difficulty of continuous casting production increases, the problem of grain boundary oxidation on the steel sheet surface occurs, and the manufacturing cost of the steel sheet increases.

[0006] The specification of Chinese Patent No. 106574348 specifically controls the ratio of the content of four elements such as C, Mn, Si, and Cr in alloy design. By adding 0.40 - 3% of Mn, a sufficiently low martensite phase transformation start temperature Ms can be obtained during the cooling period of the hot stamping process, realizing an improvement in tensile strength and an improvement in anti-delayed fracture property. Also, this patent shows that when Ni is concentrated on the surface of the sheet or component in a specific form, 0.25 - 2% of Ni can significantly reduce the sensitivity to delayed fracture. However, Ni is an expensive alloy element, and thus, it is well known that the addition of Ni does not contribute to cost-effective hot stamping of the steel sheet and its forming components. Also, in this patent, although the tensile strength increases due to a low Ms temperature, in actual production, a hard and brittle twin martensite is formed due to a low Ms temperature, which does not lead to an improvement in the toughness of the hot stamping processed component.

[0007] These conventional techniques fail to properly control the martensite phase transformation characteristics, resulting in the formation of hard and brittle twinned martensite, which not only interferes with subsequent processing, but also significantly affects the toughness and delayed cracking resistance of the hot stamping steel and its formed components.

[0008] In summary, to meet the demand for lightweight vehicle bodies, it is necessary to develop hot stamping steel components with a strength of 1700 MPa or more, sufficient toughness and resistance to delayed cracking, hot stamping steel sheets used in their manufacture, and manufacturing methods for the same. [Prior art documents] [Patent documents]

[0009] [Patent Document 1] Chinese Patent No. 106399837 [Patent Document 2] Chinese Patent No. 108374127 [Patent Document 3] Chinese Patent No. 106574348 Summary of the Invention [Problem to be solved by the invention]

[0010] The present invention has been made in view of the above problems existing in the prior art. [Means for solving the problem]

[0011] An object of the present invention is to provide a steel sheet for hot stamping having improved martensitic transformation properties, wherein the steel sheet substrate of the steel sheet for hot stamping of the present invention contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95%, and inevitable impurities, and the martensitic transformation finish temperature Mf of the steel sheet substrate satisfies Mf≧230°C; and The hardenability coefficient h of the steel plate substrate is calculated by the following formula and satisfies h≧8.0: h=[6.9[Mn] 2 +3.2[Si]+22.6[Cr] 2 +23.1[Mo]+(13.0[Cr]+2.5[Mo]+9.7)[Ni]+7.9[B]×10 3 +2.5] × [C]; where [X] is the mass percent of the alloying element.

[0012] Mf is measured by experiment, and the test method is described in detail in the specific embodiments, referring to "YB / T 5127-2018 Critical Point Measurement of Steel".

[0013] The formula for calculating h is h = [6.9[Mn]] given in "Progress in Research on Prediction Model of Hardenability of Steel" by Zhang Guoqiang et al. (Acta Metallurgica Sinica, Vol. 44, No. 4, pp. 224-228, 2019). 2 +3.2[Si]+22.6[Cr] 2 +23.1[Mo]+(13.0[Cr]+2.5[Mo]+9.7)[Ni]+7.9[B]×10 3 +2.5 × [C].

[0014] Preferably, the hot stamping steel sheet substrate further contains, in mass percent, at least one of 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%.

[0015] Preferably, 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%.

[0016] Preferably, the steel plate substrate of the steel plate for hot stamping of the present invention contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the remainder being Fe and unavoidable impurities.

[0017] Preferably, the steel sheet substrate of the hot stamping steel sheet contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the following: 0.01%≦W≦0.30%, 0.01%≦Mo≦0.30%. %, at least one of 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%, and 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%, and the balance being Fe and unavoidable impurities.

[0018] Preferably, Mf≧235°C.

[0019] Preferably, 8.3≦h≦13.5, more preferably 8.4≦h≦12.5, and even more preferably 8.5≦h≦11.5.

[0020] Preferably, 0.28%≦C≦0.345%.

[0021] Preferably, 0.15%≦Al≦0.38%.

[0022] Preferably, Al / N≧65, more preferably Al / N≧75.

[0023] Preferably, Si+Cr≦0.70%.

[0024] Preferably, 0.80%≦Mn≦1.45%.

[0025] Preferably, 0.10%≦Si≦0.40%, 0.01%≦Cr≦0.40%, and Si+Cr≦0.50%.

[0026] Preferably, 0.02%≦Nb+Ti≦0.06%. Preferably, 0.11%≦V≦0.20%.

[0027] Preferably, the average ten-point Vickers hardness of the steel plate substrate does not exceed 300HV0.3.

[0028] Preferably, the microstructure of the steel sheet substrate has, by area percentage, martensite + bainite ≦30%, the remainder being ferrite + pearlite.

[0029] Preferably, an aluminum alloy coating is applied to at least one surface of the steel sheet substrate.

[0030] Another object of the present invention is to provide a hot stamped component having high strength, high toughness, and improved delayed crack resistance, the hot stamped component of the present invention comprising, from inside to outside, a steel sheet substrate and an outer layer, the steel sheet substrate containing, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95%, and inevitable impurities; wherein the Mf temperature of the steel plate substrate satisfies Mf≧230°C; and The hardenability coefficient h of the steel plate substrate is calculated by the following formula, and satisfies h≧8.0: h=[6.9[Mn] 2 +3.2[Si]+22.6[Cr] 2 +23.1[Mo]+(13.0[Cr]+2.5[Mo]+9.7)[Ni]+7.9[B]×10 3 +2.5] × [C]; where [X] is the mass percent of the alloying element.

[0031] Preferably, the steel sheet substrate of the hot stamped component further contains, in mass percent, at least one of 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%.

[0032] Preferably, 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%.

[0033] Preferably, the steel sheet substrate of the hot stamped component contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the remainder being Fe and unavoidable impurities.

[0034] Preferably, said hot stamping formed component is made from the above-mentioned steel sheet for hot stamping of the present invention.

[0035] Preferably, the outer layer is composed of a decarburized layer of 1 to 20 μm, and the hardness of the decarburized layer does not exceed 50% of the hardness at the center of the steel sheet substrate, i.e., 50% of the hardness at the center of the steel sheet substrate from the outside to the inside along the thickness direction of the hot stamped component is at the boundary between the decarburized layer and the steel sheet substrate.

[0036] Preferably, the outer layer is a coating consisting of an interdiffusion layer of 4 to 15 μm and an intermetallic compound layer of Fe and Al outside the interdiffusion layer, and the Fe content of the interdiffusion layer is 70% by mass or more. That is, the Fe content of 70% is the boundary between the interdiffusion layer and the intermetallic compound layer from the outside to the inside along the thickness direction of the hot stamped component. Preferably, the thickness of the outer layer is 5 to 40 μm.

[0037] Preferably, the microstructure of the steel sheet substrate of the hot stamped component consists, by area percentage, of less than 4% bainite, less than 3% austenite, less than 3% ferrite, the balance being dislocated martensite and less than 0.4% fine alloyed carbides and / or nitrides formed by V, Nb or Ti.

[0038] Preferably, the sum of bainite + austenite + ferrite in the microstructure of the steel substrate of the hot stamped component, in area percentage, does not exceed 5%.

[0039] Preferably, the microstructure of the steel sheet substrate of the hot stamped component contains, by area percentage, 0.05-0.3% of fine alloy carbides and / or nitrides formed by V, Nb, Ti, etc., with an average grain size of 2-30 nm.

[0040] Preferably, the hot stamping molded component has a yield strength of 1200 to 1450 MPa, a tensile strength of 1750 to 2100 MPa, an elongation of 5% or more, and a breaking strain of 0.22-(UTS-1700) / 5000 or more.

[0041] Preferably, the hot stamped component has a yield strength of 1300 to 1600 MPa, a tensile strength of 1700 to 2050 MPa, an elongation of 5% or more, and a breaking strain of 0.22 or more. Also preferably, the hot stamped component has a yield strength of 1230 to 1420 MPa, a tensile strength of 1830 to 2030 MPa, an elongation of 5% or more, and a breaking strain of 0.23-(UTS-1750) / 5000 or more.

[0042] Preferably, the hot stamping molded component has a yield strength of 1350 to 1550 MPa, a tensile strength of 1750 to 2000 MPa, an elongation of 5% or more, and a fracture strain of 0.23 or more.

[0043] Another object of the present invention is a method for manufacturing a pre-plated steel sheet for the hot stamping formed component, said method comprising the following steps: a) Annealing treatment: Before coating, the steel sheet substrate is heated to a temperature of 740 to 870°C and kept at that temperature for 30 to 300 seconds. Here, the dew point is -30 to 5°C. The steel plate substrate contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95%, and unavoidable impurities. The martensitic transformation finish temperature Mf of the steel plate substrate satisfies Mf≧230°C. The hardenability coefficient h of the steel plate substrate is calculated by the following formula and satisfies h≧8.0: h=[6.9[Mn] 2 +3.2[Si]+22.6[Cr] 2 +23.1[Mo]+(13.0[Cr]+2.5[Mo]+9.7)[Ni]+7.9[B]×10 3 +2.5] × [C]; where [X] is the mass percent of the alloying element.

[0044] Preferably, the steel plate substrate further contains, in mass percent, at least one of 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%.

[0045] Preferably, 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%.

[0046] Preferably, the steel sheet substrate comprises, by mass percent, the balance being Fe and unavoidable impurities.

[0047] b) Heat-retention treatment: The annealed steel sheet substrate is cooled to a predetermined temperature in the range of 630 to 670°C, and then kept at that temperature for a time t. (F+P) ≦t≦100 seconds, where t (F+P) is the time required for 70% ferrite and pearlite to form in the steel substrate when the steel substrate is heated to various annealing temperatures and then cooled to 630-670°C; c) Hot-dip galvanizing treatment: The steel sheet substrate kept warm in b) is immersed in a heated galvanizing solution to perform hot-dip galvanizing treatment. The galvanizing solution has a composition, in mass percent, of 9-12% Si, 2-3% Fe, and the remainder being Al or an Al alloy and unavoidable impurities. The temperature of the galvanizing solution during this treatment is maintained in the range of 630-670°C. d) after the steel sheet substrate is pulled out of the plating solution and before the plating solution on at least one surface of the steel sheet substrate solidifies, excess plating solution on at least one surface is removed by blowing with an air knife to control the coating thickness on the at least one surface; and e) Cooling the steel sheet substrate to room temperature to obtain a pre-plated steel sheet having an aluminum or aluminum alloy coating.

[0048] Preferably, t (F+P) ≦t≦60 seconds, more preferably t (F+P) ≦t≦40 seconds. Preferably, t (F+P) does not exceed 35 seconds, and more preferably, t (F+P) does not exceed 20 seconds.

[0049] Preferably, in b), the temperature of the plating solution is not higher than the temperature at which the steel sheet substrate is maintained.

[0050] Decarburization (dew point: -30 to 5°C, preferably -28 to -5°C) before hot-dip galvanization creates an initial low-carbon zone in the pre-galvanized steel sheet. During the subsequent hot stamping process, initial interfacial migration due to interdiffusion between the coating layer and the steel sheet substrate first occurs in this initial low-carbon zone. The presence of this initial low-carbon zone reduces the amount of carbon atoms in the newly formed interdiffusion layer that diffuses and concentrates in the substrate steel sheet. This suppresses the formation of brittle martensite during subsequent cooling, contributing to improved toughness of hot-stamped components.

[0051] In the present invention, by setting the C content at 0.27 to 0.35% based on a reasonable ratio of alloying elements, the hot stamping component and hot stamping steel sheet of the present invention have Mf ≥ 230°C and h ≥ 8.0, thereby obtaining products with high strength, high toughness, and improved delayed cracking resistance. By setting Mf ≥ 230°C, the formation of hard and brittle twin martensite is suppressed during the phase transformation process, improving the toughness and delayed cracking resistance of martensite. At the same time, by setting h ≥ 8.0, the hot stamping steel sheet can be obtained with sufficient hardenability to produce at least 95%, preferably more than 97%, of martensite during hot stamping, thereby ensuring high strength. The combination of these two properties allows the hot stamping component to maintain high strength while improving toughness and delayed cracking resistance.

[0052] Specifically, the present invention aims to obtain a hot stamping component having a martensitic matrix with high strength, good toughness, and delayed cracking resistance. Martensite in steel is generally classified into two types: dislocation-type and twin-type. Dislocation martensite, also known as low-carbon lath martensite, has a high formation temperature. Due to its high dislocation density, carbon atoms are primarily concentrated on dislocation lines, forming dislocation gas clusters (dislocation complexes), rather than forming a supersaturated interstitial solid solution. This results in no significant strain in the martensite lattice, resulting in a body-centered cubic structure. Therefore, its high strength is primarily achieved by the interaction between carbon atoms and the high-density dislocations. On the other hand, because dislocations can glide during the deformation process, dislocation-type martensite combines high strength with high ductility and toughness. Twin-type martensite, also known as high-carbon lamellar martensite, has a low formation temperature. A large number of carbon atoms in the lattice exist as supersaturated interstitial solid solution, resulting in significant lattice distortion and the formation of a body-centered cubic structure and numerous fine twin substructures. The supersaturated carbon and fine twinning result in twinned martensite with ultrahigh strength. However, the fine twinning hinders dislocation glide and also causes significant lattice distortion, leading to the embrittlement of twinned martensite. Based on these considerations, this invention proposes that, in order to obtain high-strength, high-toughness martensitic steel, the carbon content and other alloying element contents of the material, combined with the characteristics of the hot stamping manufacturing process, be controlled to adjust the martensitic phase transition temperature while ensuring hardenability, thereby suppressing the formation of brittle twinned martensite in the structure. Therefore, when determining the alloying ratios, it is necessary to consider both high strength and high toughness, rather than focusing solely on high strength or high toughness.

[0053] To suppress the formation of twin martensite, the present invention proposes controlling the strength of austenite. This is because the transformation from austenite to martensite is primarily a shear phase transformation, occurring through the migration of martensite to the austenite interface. The resistance to this interface migration is primarily due to the strength of austenite. The higher the strength of austenite, the more difficult the phase interface migration becomes, making it more difficult for dislocations to glide. Therefore, the interface migration must be realized in the form of a twin substructure, which results in the formation of twin martensite. Therefore, the present invention recognizes that the strength of austenite must be reduced to suppress the formation of twin martensite. In the present invention, the strength of austenite is reduced primarily in the following two ways.

[0054] First, the C content is controlled. As is well known, C, a solid solution strengthening element in steel, strengthens austenite, facilitates the formation of twin martensite, and reduces the toughness of the final steel sheet structure. Therefore, in consideration of achieving both strength and toughness, the C content is controlled to a range of 0.27 to 0.35% in the present invention.

[0055] Next, the martensitic transformation finish temperature Mf is controlled. In the present invention, it was found that the martensitic transformation temperature of steel directly affects the formation of twin martensite, and that the martensitic transformation finish temperature Mf has a greater effect on the formation of twin martensite than the martensitic transformation start temperature. This is because, although dislocation martensite begins to form during martensitic transformation, as the temperature decreases, the generated martensite compresses the original austenite, increasing the austenite strength and making dislocation slip less likely to occur, thereby initiating the formation of twin martensite. The lower the martensitic transformation finish temperature Mf, the higher the strength of untransformed austenite at low temperatures, making it easier for twin martensite to form. In other words, for the same C content, the lower the Mf temperature, the more likely twin martensite is to form at low temperatures. Therefore, the present invention proposes controlling Mf to suppress the formation of twin martensite. In hot stamping steels and hot stamp-formed components, high-hardenability elements such as Mn and Cr are added to steel sheets in large amounts to improve hardenability and increase the martensite structure in the steel sheet, thereby increasing its strength. However, as mentioned above, adding high amounts of alloying elements lowers the Mf temperature, making it more likely that hard and brittle twin martensite will form during the phase transformation process, reducing the toughness of the component and its resistance to delayed cracking. On the other hand, excessively improving hardenability due to high alloying elements increases the difficulty of manufacturing and processing the steel sheet, making it difficult to control production quality and costs. Therefore, the present invention optimizes the content of elements such as Mn and Cr based on a C content of 0.27 to 0.35% and adds an appropriate amount of elements that increase the Mf temperature, such as Al, to achieve a higher Mf (Mf ≥ 230°C) and an appropriate hardenability coefficient h (h ≥ 8.0). This ensures strength while optimizing the toughness of the martensite itself by maximizing the dislocation martensite structure. The present inventors have found that the synergistic effect of an appropriate C content, hardenability, and a high Mf temperature effectively controls the strength of austenite during the cooling process, and during martensitic phase transformation, produces as much tough dislocation martensite as possible with a high dislocation density, while suppressing the production of brittle twin martensite.The final material therefore has high strength, high toughness, and resistance to delayed cracking.

[0056] Preferably, further limiting the upper limit of the hardenability coefficient (h≦13.5) can improve manufacturability and enhance economic efficiency. Specifically, setting h≦13.5 can improve the economic efficiency of material production. This can prevent excessive production of hard and brittle martensite during the manufacturing process (continuous casting, hot rolling, cold rolling, coating, etc.). This can maintain uniform and appropriate hardness in ingots, hot coils, and coated finished coils, reducing problems such as slab cracking, cracking during cold rolling, or steel sheet breakage. Furthermore, the shape of the steel sheet after cold rolling and coating is improved, reducing the difficulty of post-processing (e.g., trimming, punching, leveling) of annealed or coated finished products. Therefore, further setting the upper limit of hardenability can ensure that the martensite fraction after hot stamping is 95% or more, ensuring sufficient workability of the steel sheet before hot stamping. In addition, hot-rolled steel sheets and coated / annealed steel sheets in the steel sheet production link have low tensile strength, which reduces production costs.

[0057] Furthermore, the appropriate amount of Al added to the alloy and the Al / N ratio are also important considerations in the present invention. Adding a certain amount of Al not only ensures a higher Mf temperature, but also forms AlN with N to fix N, suppressing the bonding between N and B and ensuring the quenching-promoting effect of B. Furthermore, because AlN is formed at a relatively low temperature, the grain size of the formed AlN is relatively small and does not grow easily, making it less likely to seriously affect the toughness of the steel. However, adding too much Al significantly increases the austenite transformation finish temperature of the steel, potentially preventing complete austenitization during hot stamping heating. At the same time, this can also cause problems such as increased resistance at the crystallization furnace throat during continuous casting. Furthermore, adding too much Al can form grain boundary oxides and aggregates on the surface of hot-rolled steel sheets, potentially affecting the pickling effect. Therefore, in the present invention, it is necessary to control the Al content within the range of 0.11 to 0.39%. Furthermore, it was discovered that, when the N content does not exceed 0.006%, by setting the Al / N ratio to 65 or more, the effect of B in improving hardenability due to the addition of Al can be maintained, while the increased difficulty in manufacturing continuously cast links and the problem of grain boundary oxidation on the steel sheet surface can be effectively avoided.

[0058] The hot stamped component according to the present invention has sufficient toughness and delayed cracking resistance, high strength, and excellent potential for weight reduction. The hot stamped component has the following properties: yield strength of 1200 to 1450 MPa, tensile strength UTS of 1750 to 2100 MPa, elongation of ≥ 5%, and fracture strain of ≥ 0.22 - (UTS - 1700) / 5000. The high fracture strain indicates that the component has sufficient toughness and delayed cracking resistance, significantly reducing the risk of delayed cracking during storage, processing, transportation, and welding of the component after hot stamping or loading.

[0059] Those skilled in the art will appreciate that any range or any value within each of the intervals above is applicable to the present invention.

[0060] Embodiments, features and advantages of the present invention will become apparent from the following description taken in conjunction with the accompanying drawings. [Brief explanation of the drawings]

[0061] [Figure 1] FIG. 1 is a schematic diagram of the phase change expansion curve and Mf temperature calculation in the cooling process of T1 pre-plated steel sheet. [Figure 2] FIG. 1 is a schematic diagram of the phase change expansion curve and t(F+P) calculation of a T1 pre-plated steel sheet during the heat-holding process at an annealing temperature of 820°C and a heat-holding temperature of 647°C. [Figure 3] SEM microstructure of the steel substrate of T1 pre-plated steel sheet. [Figure 4] SEM microstructure of the steel substrate of CT1 pre-plated steel sheet. [Figure 5] SEM microstructure of the steel substrate of the T1 hot stamping formed component without tempering. [Figure 6] TEM microstructure of the steel substrate of the T1 hot stamping component without tempering. [Figure 7] The relationship between the martensitic phase transformation finish temperature Mf and fracture strain for hot stamped components without tempering is shown. [Figure 8] TEM microstructure of the steel substrate of the CT1 hot stamping component without tempering. [Figure 9] SEM microstructure of the steel substrate of the T5 hot stamping component without tempering. [Figure 10] 1 shows the relationship between tensile strength and fracture strain of a hot stamped component without tempering. DETAILED DESCRIPTION OF THE INVENTION

[0062] The present invention will be described in more detail below in conjunction with exemplary embodiments. The content (%) of chemical elements in this specification means weight percent. Unless otherwise specified, preferred embodiments can be freely combined as needed. Unless expressly stated, all ranges include the lower limit. Those skilled in the art will understand that the data and various parameters described in the embodiments are merely exemplary and do not limit the present invention.

[0063] The steel sheet substrate of the first exemplary hot stamped component of the present invention contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95% and unavoidable impurities, and the hardenability factor of the steel sheet substrate is h≧8.0 and Mf≧230°C.

[0064] Preferably, the steel sheet substrate of the second exemplary hot stamped component of the present invention contains, in mass percent, 0.28%≦C≦0.345%, 0.90%≦Mn≦1.45%, 0.20%≦Si≦0.50%, 0.05%≦Cr≦0.50%, 0.0017%≦B≦0.005%, 0.15%≦Al≦0.38%, 0.002%≦N≦0.0045%, 0.02%≦Nb+Ti≦0.06%, 0.14%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95% and unavoidable impurities, and the hardenability factor of the steel sheet substrate is h≧8.0 and Mf≧230°C.

[0065] Preferably, the steel sheet substrate of the third exemplary hot stamped component of the present invention contains, in mass percent, 0.28%≦C≦0.31%, 1.05%≦Mn≦1.40%, 0.20%≦Si≦0.35%, 0.06%≦Cr≦0.15%, 0.0018%≦B≦0.004%, 0.20%≦Al≦0.36%, N≦0.004%, 0.002%≦Nb+Ti≦0.05%, 0.10%≦V≦0.16%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95% and unavoidable impurities, and the hardenability coefficient h of the steel sheet substrate is ≧8.0 and Mf≧230°C.

[0066] Preferably, the steel sheet substrate of the first, second and third exemplary hot stamped components has, in mass percent, the balance Fe and unavoidable impurities.

[0067] More preferably, the steel sheet substrate of the first, second and third exemplary hot stamping formed components further includes, in mass percent, at least one of 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%.

[0068] Preferably, 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%.

[0069] Preferably, the steel sheet substrate of the hot stamped component of the present invention contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the following: 0.01%≦W≦0.30%, 0.01%≦Mo≦0.30%, 0.01%≦Ni≦0.100%. and at least one of 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦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%, and 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%, with the balance being Fe and unavoidable impurities, and the hardenability coefficient of the steel plate base material is h≧8.0 and Mf≧230°C.

[0070] Preferably, the steel sheet substrate of the hot stamped component of the present invention contains, in mass percent, 0.28%≦C≦0.345%, 0.90%≦Mn≦1.45%, 0.20%≦Si≦0.50%, 0.05%≦Cr≦0.50%, 0.0017%≦B≦0.005%, 0.15%≦Al≦0.38%, 0.002%≦N≦0.0045%, 0.02%≦Nb+Ti≦0.06%, 0.14%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the following: 0.01%≦W≦0.30%, 0.01%≦Mo≦0.30%, 0.01%≦W≦0.30%, 0.01%≦Mo≦0.30%, 0.01%≦Ti≦0.06%, 0.14%≦V≦0.20%, 0.01%≦P≦0.100%, 0.01%≦S≦0.100%. and 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%, the balance being Fe and unavoidable impurities, and the hardenability coefficient of the steel plate base material is h≧8.0 and Mf≧230°C.

[0071] Preferably, the steel sheet substrate of the hot stamped component of the present invention contains, in mass percent, 0.28%≦C≦0.31%, 1.05%≦Mn≦1.40%, 0.20%≦Si≦0.35%, 0.06%≦Cr≦0.15%, 0.0018%≦B≦0.004%, 0.20%≦Al≦0.36%, N≦0.004%, 0.002%≦Nb+Ti≦0.05%, 0.10%≦V≦0.16%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, and the following: 0.01%≦W≦0.30%, 0.01%≦Mo≦0.30%, 0.01%≦ The steel sheet substrate contains at least one of 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%, and 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%, with the balance being Fe and unavoidable impurities, and the hardenability coefficient of the steel sheet substrate is h≧8.0 and Mf≧230°C.

[0072] The chemical composition of the steel sheet substrate of the hot stamped component of the present invention is as follows: C: 0.27 to 0.35% Carbon is the most effective alloying element for improving the strength of steel sheets. In martensitic steels, carbon is an important interstitial solid-solution strengthening element, exhibiting a significant solid-solution strengthening effect. Therefore, a carbon content of less than 0.27% does not sufficiently increase the strength of martensite. However, excessive carbon content can lead to the formation of brittle twinned martensite. Therefore, to ensure the toughness of martensite, it is necessary to avoid an excessively high carbon content. A high carbon content also reduces the weldability of the material. Furthermore, the carbon content in steel significantly affects the phase change characteristics of the steel, particularly the metal-fiber temperature (Mf). The higher the carbon content, the lower the metal-fiber temperature, and at low Mf temperatures, a hard and brittle twinned martensite structure is more likely to form. Therefore, to ensure high strength, good toughness, and weldability of components, the carbon content in this invention is set to a range of 0.27 to 0.35%. In order to ensure a better balance between strength and toughness of the steel material, the C content is preferably 0.28 to 0.345%, and more preferably 0.28 to 0.31%.

[0073] Mn: 0.70 to 1.55% Mn improves hardenability and significantly enhances austenite stability. Adding a certain amount of Mn to an alloy composition can lower the austenite transformation finish temperature during the heating process, thereby expanding the temperature range for full austenitization. It can also delay the formation of ferrite and bainite during the cooling process, accelerating martensite phase transformation. For this reason, a certain amount of Mn is typically added to hot-stamped components. However, if the Mn content exceeds 1.55%, the strength improvement effect is small, segregation control becomes difficult, and the Mf temperature drops significantly, significantly reducing the toughness and delayed cracking resistance of the hot-stamped component. Therefore, in the present invention, the Mn content is limited to a range of 0.70 to 1.55%. Preferably, the Mn content is in the range of 0.80 to 1.45%.

[0074] N≦0.006% N is an impurity element in steel. In particular, in steels containing B, when N combines with B, the hardenability-improving effect of B is significantly reduced. Therefore, the N content must be kept as low as possible. However, in industrial production, various processes for removing N all increase production costs. Therefore, taking into consideration the overall manufacturing cost, the N content in the present invention is set to 0.006% or less.

[0075] 0.11%≦Al≦0.39%, Al / N≧65 Al is an important element in this invention. First, Al is a strong deoxidizing element and is often used as a deoxidizer in steel smelting. Second, Al is a ferrite stabilizer, increasing the Mf temperature. This effect is the opposite of the effect of elements such as C, Mn, and Cr on the Mf temperature, and helps suppress the formation of hard and brittle twin martensite. Furthermore, Al forms N and AlN and solidifies N, thereby suppressing the bonding between N and B and ensuring the hardenability-improving effect of B. Furthermore, because the formation temperature of AlN is relatively low, the resulting AlN particles are relatively small, which is unlikely to seriously affect the toughness of the steel. For this reason, an Al content of 0.11% or more is desirable. However, excessive Al content significantly increases the austenite transformation finish temperature of the steel, preventing complete austenitization during hot stamping. Furthermore, problems such as increased resistance at the crystallization outlet during continuous casting and difficulty in pickling due to grain boundary oxidation on the surface of hot-rolled steel sheets can occur. Therefore, in the present invention, the Al content is set to a range of 0.11 to 0.39%, preferably 0.15 to 0.38%. Furthermore, the present invention has found that by limiting Al / N to 65 or more, the hardenability improving effect of B is ensured, while avoiding the increase in the difficulty of manufacturing continuously cast links and the problems of grain boundary oxidation on the steel sheet surface, and the content range of the Al element can be further controlled.

[0076] Si: 0.10 to 0.40% Silicon dissolves in the base metal and has the effect of improving the strength of the base metal. It can also be used as a deoxidizer in the steelmaking process. Therefore, it is necessary to add more silicon. However, if the silicon content is too high, the degree of grain boundary oxidation and decarburization varies depending on the position on the steel sheet surface during the cooling process of the steel sheet coil after hot rolling, which affects the pickling effect of the steel sheet coil and the surface quality of the final product. Therefore, in the present invention, the silicon content is set to a range of 0.10 to 0.40%.

[0077] Cr: 0.01 to 0.40% Cr is an element that improves the hardenability of steel. At the same time, it is highly effective in preventing oxidation and surface decarburization of steel. However, if the Cr content is too high, pickling of the oxide scale on the surface of the hot-rolled coil becomes difficult, affecting the surface quality of the final product. Therefore, in the present invention, the Cr content is set to the range of 0.01 to 0.40%.

[0078] Si+Cr≦0.70% Si and Cr reduce the pickling properties of hot-rolled coils through oxidation, leading to a deterioration in the surface quality of the final product. Therefore, in the present invention, the total content of Si and Cr is preferably 0.70% or less, and more preferably 0.50% or less.

[0079] B: 0.001 to 0.01% B segregates at austenite grain boundaries during hot stamping, suppressing the formation of ferrite and significantly improving the hardenability of steel. For this reason, a certain amount of B is added in the present invention. However, too much B content causes boron embrittlement, adversely affecting performance improvement. Therefore, in the present invention, the B content is set to the range of 0.001 to 0.01%.

[0080] 0.001%≦Nb+Ti≦0.1% and 0.05%≦V≦0.20% V, Nb, and Ti form carbides in steel, which play an important role in grain refinement and precipitation strengthening. At the same time, the formation of these carbides consumes the carbon content in the matrix, further improving the strength and toughness of the final product. However, excessive addition of these elements not only fails to further improve strength and toughness, but also significantly increases production costs. Furthermore, Ti and N have strong bonding strength and form TiN in solid solution with N in steel. Although the bonding between N and B is prevented and the effect of B is maintained, TiN precipitates at high temperatures and tends to grow into coarse TiN particles during the subsequent slab cooling process, thereby not contributing to improving the toughness of the steel sheet. Therefore, the total content of Nb and Ti in the present invention is 0.001 to 0.10%, and the V content is 0.05 to 0.20%. Preferably, the total content of Nb and Ti is 0.02 to 0.06%, and the V content is 0.11 to 0.20%.

[0081] P: 0.001% to 0.100% P is an element that is inevitably contained in steel. As a solid solution strengthening element, P can improve the strength of steel sheet relatively inexpensively. On the other hand, if the P content exceeds 0.100%, P segregates at grain boundaries, causing adverse effects such as a significant decrease in toughness. Therefore, the upper limit of the P content is set to 0.100% or less, preferably 0.050% or less. Furthermore, considering that a P content of less than 0.001% increases smelting costs, the lower limit of the P content is set to 0.001% or more, preferably 0.004% or more.

[0082] S: 0.0001% to 0.100% Like P, S is an element that is inevitably contained in steel and reacts with Mn in the steel to form inclusions in the steel in the form of MnS. If the S content exceeds 0.100%, the large amount of MnS significantly impairs the ductility and toughness of the steel, reducing workability. Therefore, the upper limit of the S content is set to 0.100% or less, preferably 0.015% or less. Similarly, considering that reducing the S content to less than 0.0001% increases smelting costs, the lower limit of the S content is set to 0.0001% or more, preferably 0.0005% or more, and more preferably 0.001% or more.

[0083] Furthermore, other elements that may be contained in the steel sheet substrate of the hot stamped component will be described.

[0084] W, Mo, Ni, Cu, Co:0.01~0.30% The addition of W, Mo, Ni, Cu, and Co can improve the hardenability of steel, but because the addition of these alloying elements increases the alloy cost, only appropriate amounts are added to the material. At the same time, Ni, Cu, and Co also have the advantage of improving the toughness of the material. When the content of each element is 0.01% or more, the above beneficial effects can be achieved. Therefore, it is preferable to set the lower limit of the content of each of these five elements to 0.01% or more. At the same time, considering the economic viability of the material alloy, it is preferable to set the upper limit of the content of each of these five elements to 0.30% or less. In this case, toughness can be improved, the impact on the hardenability of the steel is reduced, and the machinability of the steel is ensured.

[0085] Sn:0.005~0.30%, Sb:0.005~0.100% When the Sn and Sb contents are 0.005% or more, the wettability of the coating can be improved. Therefore, the lower limits of the Sn and Sb contents are preferably 0.005% or more. However, if the Sn content exceeds 0.300% or the Sb content exceeds 0.100%, the toughness of the material decreases, so it is preferable to set the Sn content to 0.300% or less and the Sb content to 0.100% or less.

[0086] Ca:0.0001~0.01%, Mg:0.0001~0.01%, Zr:0.0001~0.01%, REM:0.0001~0.01% The inclusion of 0.0001% or more of REM (rare earth metal) can contribute to purifying molten steel, controlling the morphology and distribution of impurities in steel, and refining crystal grains. Ca, Mg, and Zr each have the effect of refining inclusions at a content of 0.0001% or more, leading to improved material performance. Therefore, the contents of Ca, Mg, Zr, and REM are preferably 0.0001% or more. On the other hand, when the content of each element exceeds 0.01%, the above effects saturate. Therefore, the contents of Ca, Mg, Zr, and REM are preferably 0.01% or less.

[0087] Preferably, in consideration of the cost of the alloy, the total content of W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM elements is in the range of 0.0001% to 0.30%. Generally, the steel sheet substrate of the present invention may contain elements such as W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM, but the presence of these elements does not affect the solution of the technical problem of the present invention.

[0088] In this embodiment, there are no particular limitations on other components of the steel sheet substrate of the hot stamping component. For example, scrap may contain elements such as As, but as long as these elements are within normal ranges, they do not affect the properties of the steel sheet substrate of the hot stamping steel sheet.

[0089] To achieve high-strength components, when the C content is high, it is desirable for the steel sheet to obtain as much martensite structure as possible after hot stamping, which requires the steel sheet to have sufficient hardenability. Numerous experiments have shown that the hardenability factor h of the component of the present invention should be 8.0 or higher to ensure the formation of sufficient martensite phase (95% or more by area percentage) and obtain a tensile strength of 1700 MPa or more. At the same time, in order to improve the toughness and delayed cracking resistance of the formed hot-stamped component, the present invention also proposes that the component have a high Mf temperature. Specific explanations are as follows.

[0090] Generally, increasing the content of alloying elements such as C, Mn, and Cr improves the strength of a material. Because these alloying elements are all stable austenitizing elements, austenite can exist at lower temperatures, thereby improving the strength of austenite. Because martensitic phase transformation is a shear mechanism, dislocation slip is likely to occur when the austenite strength is relatively low. Therefore, martensitic phase transformation is primarily based on dislocation deformation, resulting in the formation of dislocation-type lath martensite. When the austenite strength is relatively high, martensitic phase transformation is primarily based on twinning deformation, resulting in the formation of twinned martensite. Therefore, if not properly controlled, high-alloying elements such as C, Mn, and Cr in steel can lower the end temperature of martensitic phase transformation, improving austenite strength and increasing the likelihood of the formation of hard, brittle twinned martensite, which can impair the toughness of the steel. This is also one of the important reasons why hot-stamped components with tensile strengths of 1700 MPa or more are prone to delayed cracking during storage, transportation, and subsequent welding after hot stamping. Therefore, the present invention emphasizes that when developing a hot-stamped component with a strength exceeding 1700 MPa, it is necessary to achieve high strength while improving toughness and delayed cracking resistance by appropriately blending other alloying elements, while maintaining the appropriate C content, to simultaneously satisfy the requirements for hardenability and Mf temperature. Specifically, while maintaining hardenability, the contents of elements such as Mn and Cr are adjusted to reduce the effect of lowering the martensitic transformation finish temperature, and an appropriate amount of alloying elements such as Al that raise the Mf temperature is added to achieve a sufficiently high Mf temperature. In the present invention, it was discovered that by setting the C content to 0.27-0.35%, Mf ≥ 230°C, and hardenability h ≥ 8.0, martensitic transformation occurs at a higher temperature, appropriately controlling austenite strength. This allows the martensitic transformation to produce as much dislocation martensite as possible, suppressing the formation of brittle twin martensite. As a result, the final material has high strength, high toughness, and improved delayed cracking resistance. Furthermore, when the Mf temperature exceeds 235°C, the toughness and delayed cracking problems of the steel plate are further improved.

[0091] Preferably, to improve the manufacturability of hot stamping steel sheets, the hardenability of the steel sheets should not be too high. This prevents excessive production of hard and brittle martensite during the manufacturing process (continuous casting, hot rolling, coating, etc.) of the steel sheets, ensuring uniform and appropriate hardness in the cast billets, hot coils, and coated finished coils. This reduces the occurrence of problems such as slab cracking, cold rolling cracking, and steel sheet breakage, and improves sheet shape during the cold rolling and coating processes. At the same time, it also reduces the difficulty of subsequent processes (trimming, punching, leveling, etc.) of the pre-coated steel sheets. Therefore, after numerous experimental verifications, it has been determined that the hardenability factor h of the components of the present invention should be 13.5 or less, preferably 12.5 or less, or even 11.5 or less, so that the resulting hot stamping components can achieve both high mechanical properties and improved manufacturability.

[0092] A hot stamped component is composed of a steel sheet substrate and an outer layer from the inside to the outside. If a pre-plated layer is not present on the surface of a steel sheet for hot stamping, a decarburized layer of a certain thickness is formed on the surface during the hot stamping process, and this decarburized layer becomes the aforementioned outer layer. The presence of a decarburized layer of a certain thickness allows the component to have good bending toughness. However, as the decarburized layer becomes thicker, not only does the component strength decrease, but also the peak load in a VDA bending test decreases. Therefore, in the present invention, the thickness of the decarburized layer is set to 1 to 20 μm and is defined based on the distance from the position where the hardness reaches 50% of the hardness of the center of the steel sheet substrate in a Vickers hardness test to the surface of the hot stamped component. If a pre-plated layer of aluminum or aluminum alloy is present on the surface of a steel sheet for hot stamping, elements in the aluminum or aluminum alloy pre-plated layer and the steel sheet substrate diffuse during the hot stamping process, forming an interdiffusion layer and an intermetallic compound layer of Fe and Al. In this case, the outer layer of the hot-stamped component is a coating composed of an interdiffusion layer and an intermetallic compound layer of Fe and Al outside the interdiffusion layer, and the thickness of the coating is 5 to 40 μm. The thickness of the interdiffusion layer is 4 to 15 μm. The interdiffusion layer is mainly α-Fe containing a large amount of Al and Si, and the Fe content thereof is 70 wt % or more.

[0093] After designing the alloy, the microstructure of the steel plate substrate of the component has a composition, by area percentage, of less than 5% bainite, less than 3% austenite, and less than 3% ferrite, with the remainder being dislocation martensite, and less than 0.4% carbides and / or nitrides formed by V, Nb, Ti, etc. Preferably, to further improve strength and toughness, the less non-martensitic phases in the microstructure, the better, and the total of bainite, ferrite, and retained austenite should be 5% or less. Furthermore, fine alloy carbides and / or nitrides formed by minutely dispersed elements such as V, Nb, and Ti have effects such as precipitation strengthening, grain refinement, and improving martensite toughness by consuming C in martensite. To achieve these effects, the average particle size of the fine alloy carbides and / or nitrides formed by V, Nb, Ti, etc. should be controlled to 2-30 nm, and their content should be in the range of 0.05-0.4%.

[0094] Based on the above alloy design and microstructural characteristics, the hot stamped component of the present invention exhibits ultra-high strength, good plasticity, and toughness. Its yield strength is 1200-1450 MPa, its tensile strength UTS is 1750-2100 MPa, its elongation is 5% or more, and its fracture strain is 0.22-(UTS-1700) / 5000 or more. Preferably, its yield strength is 1230-1420 MPa, its tensile strength UTS is 1830-2030 MPa, and its fracture strain is 0.23-(UTS-1750) / 5000 or more. After tempering, the hot stamped component of the present invention has a yield strength of 1300-1600 MPa, a tensile strength UTS of 1700-2050 MPa, an elongation of 5% or more, and a fracture strain of 0.22 or more. Preferably, the yield strength is 1350 to 1550 MPa, the tensile strength is 1750 to 2000 MPa, and the breaking strain is 0.23 or more. In the tempering treatment, the steel is kept at a temperature of 150 to 230°C for 30 to 80 minutes, and then cooled outside the furnace.

[0095] The present invention also provides a hot stamping steel sheet for producing the above-mentioned hot stamped component, wherein the steel sheet substrate of the hot stamping steel sheet contains, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0.100%, Fe≧95%, and unavoidable impurities, and the Mf of the steel sheet substrate satisfies Mf≧230°C, and the hardenability coefficient h satisfies h≧8.0.

[0096] Preferably, the steel plate substrate of the hot stamping steel plate further contains, in mass percent, at least one of the following: 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%.

[0097] Preferably, 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0.30%.

[0098] Preferably, the steel sheet substrate of the steel sheet for hot stamping has, in mass percent, the balance being Fe and unavoidable impurities.

[0099] Preferably, the steel sheet for hot stamping is a pre-plated steel sheet, i.e., an aluminum or aluminum alloy coating is applied to the outside of a steel sheet substrate. To reduce the strength of the finished product and improve its formability, the microstructure of the steel sheet after coating and cooling to room temperature is primarily ferrite and pearlite. The formation of martensite should be avoided as much as possible, with its proportion being less than 30%, preferably 10% or less. As mentioned above, lowering the upper limit of h can prevent the formation of martensite during processing of the steel sheet as much as possible. The organization of ferrite and pearlite reduces the hardness of the steel sheet, with the average 10-point Vickers hardness not exceeding 300 HV0.3, preferably not exceeding 260 HV0.3. Preferably, the thickness of the aluminum or aluminum alloy pre-plated layer is 5 to 20 μm.

[0100] The method for producing the pre-plated steel sheet includes the following steps: a) Before coating, a steel sheet substrate having the above composition is heated to 740-870°C in an H2 and N2 atmosphere (H2 volume ratio 2-12%) and annealed by holding for 30-300 seconds to a dew point of -30-5°C. If the dew point is too high, oxidation of the steel sheet surface will progress significantly, affecting the coating quality. On the other hand, if the dew point is too low, the desired surface decarburization effect will not be achieved. b) After annealing, the steel sheet substrate is cooled to a predetermined temperature within the range of 630 to 670°C and kept at that temperature for a certain time t. (F+P) ≦t≦100 seconds, where t (F+P) is the time required to heat the steel sheet substrate with the above composition to various annealing temperatures, then rapidly cool it to 630-670°C and hold it at that temperature until 70% of ferrite and pearlite are formed; Typically, the length of each stage in a coating and plating production line is fixed, and the residence time of the steel sheet at each stage can be adjusted by adjusting the moving speed of the steel sheet. Meanwhile, to ensure that the steel sheet has a structural characteristic consisting primarily of ferrite and pearlite, with a martensite content of less than 30%, the steel sheet must be kept at a sufficient temperature before immersion in the plating solution to allow the transformation of ferrite and pearlite to occur. The holding time t is the time t required to produce 70% ferrite and pearlite after heating the steel sheet substrate to various annealing temperatures, rapidly cooling to 630-670°C, and then holding the temperature. (F+P) On the other hand, if the moving speed of the steel sheet is too slow, production efficiency will drop significantly and manufacturing costs will increase. Therefore, the holding time t must be 100 seconds or less, preferably 60 seconds or less, and more preferably 40 seconds or less.

[0101] c) The steel sheet substrate after the warming step b) is immersed in a heated coating solution for hot dip coating. Here, the composition of the coating solution is, in mass percent, 9-12% Si, 2-3% Fe, and the balance being Al or Al alloy and unavoidable impurities. In this process, the temperature of the coating solution is maintained at 630-670°C. To prevent a decrease in the temperature of the coating solution after immersion of the steel sheet in the coating solution, which would increase the amount of Fe slag and affect the quality of the pre-coating, the temperature of the coating solution should preferably not be higher than the temperature used in the warm-drying step b). d) After the steel sheet substrate is pulled out of the plating solution, before the plating solution on at least one surface of the steel sheet substrate solidifies, excess plating solution on at least one surface is removed by air knife blowing to control the thickness of the pre-plating layer on at least one surface; and e) The coated steel sheet substrate is cooled to room temperature to obtain a pre-plated steel sheet having an aluminum or aluminum alloy pre-plated layer.

[0102] The manufacturing method of the present invention improves production efficiency while reducing the occurrence of problems such as slab cracking, cold rolling cracking, and steel plate breakage by making the structural characteristics of the steel plate substrate primarily ferrite and pearlite. Furthermore, the plate shape of the steel plate after cold rolling and coating is improved, which reduces the difficulty of subsequent processes (trimming, punching, leveling, etc.) of annealed or coated finished products.

[0103] The present invention will be described in more detail below based on the embodiments. The following examples and experimental data are provided to exemplify the present invention, and it will be apparent to those skilled in the art that the present invention is not limited to these examples or experimental data.

[0104] Steel sheets having the compositions shown in Table 1 were prepared, and the corresponding manufacturing process was as follows. The test steel plates are manufactured by the following process. a) Steelmaking: According to the composition in Table 1, steel is melted in a vacuum induction furnace, an electric furnace or a converter, and then ingots are produced by continuous casting technology, or thin slab continuous casting and rolling process is directly adopted. b) Hot rolling: The steel ingot is heated to 1200°C and held there for 2 hours, then hot rolled at 800°C to 1200°C and coiled at 600°C to form a hot-rolled coil. The hot-rolled coil is then pickled to remove the oxide scale formed during hot rolling. c) Cold rolling: The pickled hot rolled coil is cold rolled to obtain a cold rolled steel coil with a thickness of 1.2 mm, with a cold rolling reduction of 30% to 70%.

[0105] [Table 1]

[0106] T1 to T8 are examples of the present invention, and CT1 to CT3 are comparative examples. Among these, h is calculated according to the formula in the present invention, and Mf temperature is measured using a DIL805A phase change measurement device in accordance with the "YB / T5127-2018 Steel Critical Point Measurement" standard. Specific test methods are as follows: The sample is heated to 930°C at a heating rate of 10°C / s, held at that temperature for 300 seconds, and then cooled to below 100°C at a cooling rate of 30°C / s. The longitudinal expansion deformation of the sample during the entire cooling process is recorded, and the temperature corresponding to 95% martensitic phase transformation is determined using this principle as the martensitic phase transformation finish temperature, Mf. Figure 1 illustrates the Mf temperature test for a T1 test steel plate.

[0107] The above manufacturing method for pre-plated steel sheets is carried out according to the parameters T1 to T8 and CT1 to CT3 in Table 2. The composition of the plating solution is, in mass percent, 9 to 12% Si, 2 to 3% Fe, and the remainder Al and unavoidable impurities. The temperature of the plating solution is maintained at 630 to 670°C. The Vickers hardness of the steel substrate of the pre-plated steel sheets is then tested. The indenter load is 0.3 kg·F, and the test results are the average of 10 points. Table 2 shows the hardness values ​​of the steel substrate. (F+P) was obtained by simulating the manufacturing method of the corresponding plated steel sheet using the DIL805A phase change device. Figure 2 shows the phase change expansion curves of the T1 pre-plated steel sheet during the heat-holding process with an annealing temperature of 820°C and a heat-holding temperature of 647°C. (F+P) 3 and 4 show the microstructures of the steel sheet substrates of the obtained T1 pre-plated steel sheets and CT1 pre-plated steel sheets.

[0108] [Table 2]

[0109] As shown in Figure 3, the steel substrate of T1 pre-galvanized steel sheets is mainly composed of ferrite and pearlite, with a martensite content of less than 5%. The typical microstructure of T1 is applied to the steel substrate of T2 to T8 pre-galvanized steel sheets, and the microstructure corresponds to the hardness of the steel substrate of T1 to T8 pre-galvanized steel sheets (227 to 275 HV0.3). The reason why the steel substrate is mainly composed of ferrite and pearlite is that the holding time before hot dip galvanizing is longer than the steel substrate temperature. (F+P)This is because the heating time is longer than the heating time, which provides sufficient time for the formation of large amounts of ferrite and pearlite, thereby reducing or avoiding the formation of martensite.

[0110] On the other hand, as shown in Figure 4, the steel sheet substrate of CT1 pre-coated steel sheet is mainly composed of martensite and ferrite, with the martensite content exceeding 30%. This microstructure results in a hardness of the steel sheet substrate exceeding 300 HV0.3. The same is true for CT2 to CT3 pre-coated steel sheets. This is because the high alloy design of CT1 to CT3 steel sheets improves the hardenability of the steel sheet, with an h value of 13.7 to 15.3, making it difficult for non-martensitic structures such as ferrite and pearlite to form during the annealing treatment of the coating process. Therefore, t (F+P) The time is relatively long, and the retention time is t (F+P) If the cooling time is shorter than this, the amount of ferrite and pearlite formed before hot-dip galvanization will be insufficient, making it difficult to control the cooling rate during the subsequent cooling process after hot-dip galvanization. This inevitably results in the formation of a large amount of martensite, resulting in an increase in the hardness of the steel sheet substrate, making it unsuitable for subsequent processing. In particular, CT3 has a relatively high hardness of the steel sheet substrate of the pre-plated steel sheet, even though all of its alloying elements are within the range of the present invention. This is because, when setting the alloy ratio, the hardenability h during processing of the pre-plated steel sheet is t (F+P) This is because the influence of the temperature on the formation of martensite during processing was not taken into consideration, making it impossible to effectively control the formation of martensite during processing. In addition, due to the high alloy design, the Mf temperatures of CT1 to CT3 are generally low, and the amount of hard and brittle twin martensite in the formed martensite increases, further reducing the formability of the steel sheet.

[0111] Therefore, when manufacturing pre-plated steel sheets, in combination with the martensitic phase transformation characteristics of the material, t (F+P) It is necessary to rationally set the holding time t before hot dip plating so that t≦t≦100s, ensuring production efficiency and minimizing the formation of martensite to facilitate subsequent processes. (F+P)It is preferable that t(F+P) is 35 seconds or less, and it is more preferable that t(F+P) is 2 seconds or less and t is 60 seconds or less.

[0112] Hot stamping plate simulation Cold-rolled steel plate T1 * and T4 * A hot stamping simulation was conducted using (uncoated), T1 to T8 pre-coated steel sheets, and CT1 to CT3 pre-coated steel sheets. After heating to 920°C in a heating furnace for 300 seconds, the samples were transferred to a press and hot stamped for 8 to 12 seconds. The corresponding hot stamped components were then obtained by cooling to below 200°C at a cooling rate of 40°C / s. Furthermore, to simulate the paint baking process for automotive parts, the above hot stamped components were tempered. The tempering process was set to 170°C for 20 minutes, based on actual baking processes.

[0113] Measuring the outer layer thickness of hot stamped components The hot-stamped components T1-T8 and CT1-CT3, formed by hot stamping pre-plated steel sheets, comprise, from the inside to the outside, a steel sheet substrate and an outer layer. The outer layer is a coating consisting of an interdiffusion layer and an Fe-Al intermetallic compound layer outside it. The microstructure of the steel sheet substrate for each component was observed using a metallographic microscope and a scanning electron microscope, and the coating was subjected to EDS line scanning to measure changes in Fe content, thereby measuring the overall coating thickness and the thickness of the interdiffusion layer. The thickness of the interdiffusion layer was measured at the 70% Fe content, which is the boundary between the interdiffusion layer and the Fe-Al intermetallic compound layer.

[0114] T1 is formed by hot stamping cold rolled steel sheet * and T4 *The hot stamped component includes a steel sheet substrate and an outer layer from the inside to the outside. The outer layer is a decarburized layer, and its thickness is defined as the thickness from the surface of the hot stamped component to the position where the hardness is 50% of the hardness of the center of the steel sheet substrate. The measurement method is as follows: 1) Measure the hardness of the center of the steel substrate of the hot stamped component using an indenter load of 0.3 kg·F; 2) After step 1), reduce the indenter load to 0.03 kg F and measure the hardness stepwise from the surface of the steel plate toward the inside. When the measured hardness reaches 50% of the hardness measured in step 1), record the position of the hardness measurement point and determine the distance from that point to the component surface as the thickness of the decarburized layer.

[0115] To reduce measurement errors, all results are averaged over 10 measurements.

[0116] The outer layer thickness and structural test results of all the above components are shown in Table 3.

[0117] Performance testing of hot stamped components The resulting untempered and tempered hot stamped components were tested for tensile strength, elongation, and maximum bending angle according to room temperature tensile test standard GBT 228.1 and three-point bending test standard VDA 238. To reduce measurement error, the final test results were the average of three test result groups.

[0118] The bending fracture strain test method is as follows: (1) Static three-point bending test is performed to measure the VDA bending angle α of the constituent test specimen. peak (2) Based on the experimental results, αL ≥ 50%α peak The interrupted bending angle α for the interrupted bending test is L (i.e., bending angles of the component specimens under load); (3) select at least three groups of component specimens with α L The load was stopped when the component test piece was bent to the bending angle α UL(4) Place the unloaded component specimen under an optical microscope and measure the inner radius Ri and outer radius Ro of the most severely deformed area; (5) Measure the different α according to the following equation (1): UL The equivalent (plastic) strain ε (component specimen is denoted by α) on the outer surface of the most severely deformed zone of the component specimen under the no-load condition L ε-α L (6) Based on the fitting results, the bending fracture strain ε (i.e., α L is α peak Find ε) when it is equal to .

number

[0119] The final microstructure, tensile properties, and VDA flexural properties of the hot stanned components without tempering are shown in Table 3.

[0120] [Table 3]

[0121] As shown in Table 3, the microstructures of all hot stamped components are primarily martensite (accounting for more than 95%), with less than 0.3% fine alloy carbides. In particular, the martensite content in the microstructures of the T1-T4, T6-T8, and CT1-CT3 components exceeds 97%, resulting in a nearly completely martensite structure. This nearly completely martensite structure benefits from the sufficient hardenability of each component, ensuring high strength. As an example, a typical microstructure of the T1 hot stamped component of the present invention is shown in Figures 5 and 6. It can be seen that the microstructure of the component of the present invention is dominated by dislocation martensite, which has good toughness, and no twinned martensite.

[0122] T1 to T8 and T1 of the present invention * and T4 * The components have a yield strength of 1200-1400 MPa, a tensile strength of 1850-2030 MPa, a maximum bending angle of 45-53°, and a fracture strain of 0.18-0.24, combining high strength and high toughness.

[0123] The components of the present invention achieve an Mf of 230°C or higher under conditions satisfying hardenability h≧8.0 through the appropriate alloying element ratio. For example, the Mf of each of the embodiments T1 to T8 reaches 235 to 260°C, thereby combining high strength and high toughness. Therefore, after thermal deformation, at least 95% of the martensitic structure is maintained during the die cooling process, and hard and brittle twinned martensite is not formed. The components T1 to T8 of the present invention exhibit good toughness at high strength. Figure 7 shows the relationship between Mf temperature and fracture strain for all the untempered hot stamped components obtained. It can be seen that toughness tends to increase with increasing Mf temperature. Therefore, to achieve a fracture strain of 0.18 or higher, the Mf temperature of the present invention must be 230°C or higher, and preferably, the Mf temperature is 235°C or higher.

[0124] On the other hand, the maximum bending angle of the CT1 to CT3 components did not reach 45°, and the fracture strain was less than 0.18, so they were unable to achieve both high strength and high toughness.On the other hand, the yield strength of the CT1 to CT3 components was 1200 to 1350 MPa, and the tensile strength was 1900 to 2000 MPa.

[0125] While the CT1-CT3 components achieved sufficient strength due to their high hardenability, the excessively large h values ​​and low Mf temperatures resulting from their alloy design resulted in a harder, more brittle twinned martensite structure, which did not contribute to sufficient toughness. For example, the Mf temperatures of CT1 and CT2 were only 217°C and 225°C, respectively, due to the high Mn contents of 1.79% and 1.60%, respectively (exceeding the compositional range of the present invention). Even the high Al contents of 0.59% and 0.45% (exceeding the compositional range of the present invention) were insufficient to compensate for the reduction in Mf due to Mn. Furthermore, although the composition of CT3 is within the range of the present invention, its element ratios were inappropriate, resulting in a high h of 15.28 and an Mf temperature of only 210°C. As an example, a typical microstructure of the CT1 component is shown in Figure 8. This shows twinned martensite, a result of the low Mf. Therefore, compared with the T1 to T8 components of the present invention, the CT1 component has high strength but is significantly deficient in toughness performance.

[0126] Comparing the T1 to T8 components with the CT1 to CT3 components, in order to obtain the desired toughness while maintaining high strength, the present invention requires that the alloying elements be rationally blended so that Mf≧230°C and h≧8.0, based on the composition ranges of each element according to the present invention. Furthermore, as shown in Table 3, if h≧8.0, a martensite structure of 95% or more can be obtained. In consideration of manufacturability and economic efficiency, the present invention requires that the hardenability coefficient be preferably 8.3≦h≦13.5, more preferably 8.5≦h≦11.5, and this applies to the T1 to T8 components, for example.

[0127] Furthermore, as an example, FIG. 9 shows a typical microstructure of a non-tempered T5 hot stamped component of the present invention. Comparing FIG. 5 with FIG. 9 reveals that the microstructure of the T1 component is a typical lath-type dislocation martensite structure, with a martensite fraction exceeding 97%. Meanwhile, the microstructure of the T5 component contains a small amount of pre-eutectoid ferrite. Furthermore, compared to the T1 component, the T5 component has lower strength and toughness, even with a similar composition. This is primarily due to the lower Al / N ratio of the T5 component (less than 65) compared to the T1 component. Therefore, N does not bond sufficiently with Al, and some N bonds with B, which prevents B from fully enhancing hardenability. This results in the formation of a small amount of ferrite, which does not contribute to ensuring high strength and good toughness. On the other hand, in the T1 component, a large amount of Al (0.36%) is added to make Al / N≧65, which promotes the bonding between Al and N and ensures the role of B in the steel sheet in improving hardenability, thereby preventing the formation of ferrite and preventing the formation of large TiN particles during hot stamping even in the presence of Ti, thereby further improving the strength and toughness of the component. Therefore, in order to ensure a better N fixation effect and achieve higher strength and toughness, in the present invention, it is preferable to make Al / N≧65 when 0.11%≦Al≦0.39%.

[0128] FIG. 10 shows the relationship between fracture strain and tensile strength of all the hot stamped components obtained without tempering. * and T4 * The strain at break of the component is 0.22-(UTS-1700) / 5000 or greater, and further, if the Al / N ratio is 65 or greater, the strain at break of embodiments of the present invention is 0.23-(UTS-1750) / 5000 or greater.

[0129] The final results of the tensile properties and VDA bending properties of the hot stamped components after tempering are shown in Table 4. From Table 4, T1 to T8 of the present invention, T1 * and T4 *After tempering, the components have a yield strength of 1320-1520 MPa, a tensile strength of 1780-1940 MPa, a maximum bending angle of 51-58°, and a fracture strain of 0.22-0.27, demonstrating both high strength and high toughness.

[0130] On the other hand, after tempering, the CT1 to CT3 components had a maximum bending angle of 47 to 50° and a fracture strain of 0.19 to 0.21 under conditions of a yield strength of 1350 to 1450 MPa and a tensile strength of 1800 to 1900 MPa. Although the tempering treatment relieves the internal stress of the CT1 to CT3 components and improves their toughness to some extent, compared to the embodiments of the present invention, the CT1 to CT3 components do not achieve both high strength and high toughness.

[0131] [Table 4]

[0132] Delayed cracking test The test method conformed to the four-point bending load test specified in ISO 7539-2:1995. Five samples, T1 to T8 and CT1 to CT3, of untempered hot-stamped components were taken and parallel experiments were conducted. A 0.1 mol / N hydrochloric acid solution was used. Loads of 70 to 110% of the sample's yield strength were applied. The maximum stress at which cracks did not occur after 120 hours, i.e., the ultimate stress, and its ratio to the yield strength were recorded. The results are shown in Table 5.

[0133] [Table 5]

[0134] As shown in Table 5, the T1-T8 components, even without tempering, have a 120-hour crack-free critical stress exceeding 750 MPa, and the ratio of critical stress to yield strength exceeds 55%, demonstrating excellent delayed cracking resistance. This is because the T1-T8 components have high Mf temperatures and appropriate h values ​​under conditions of hardenability. This not only ensures a nearly complete martensite structure during the cooling process after thermal deformation, but also suppresses the formation of hard, brittle twinned martensite, improving overall martensite toughness and delayed cracking resistance. In particular, the Mf values ​​for T1-T3 and T6-T8 are above 240°C, resulting in a critical stress to yield strength ratio exceeding 60%, further improving the delayed cracking resistance of the components. Furthermore, the Mf values ​​for T2 and T7 are above 255°C, resulting in a critical stress to yield strength ratio exceeding 70%, further improving the delayed cracking resistance of the components.

[0135] In contrast, the ultimate stresses of the CT1-CT3 components are relatively low, with ultimate stress to yield strength ratios of less than 50%, and delayed cracking performance is significantly inferior to the embodiments of the present invention.

[0136] In summary, the hot stamped components of the present invention exhibit excellent overall performance, providing high strength, high toughness, and improved resistance to delayed cracking.

[0137] The above embodiments and experimental data are only intended to illustrate the present invention, and it is clear to those skilled in the art that the present invention is not limited to these embodiments, and various modifications can be made without departing from the scope of protection of the present invention.

Claims

1. A steel sheet for hot stamping, wherein the steel sheet base material of the steel sheet for hot stamping is In terms of content percent, 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.1 0%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01 %, 0.11%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0 .1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001 %≦S≦0.100%, Fe≧95%, and unavoidable impurities; Here, the martensitic transformation finish temperature Mf of the steel plate substrate satisfies Mf≧230°C. ;and The hardenability coefficient h of the steel plate substrate is calculated by the following formula and satisfies h≧8.0: h=[6.9[Mn] ２ +3.2[Si]+22.6[Cr] ２ +23.1[Mo]+( 13.0 [Cr] + 2.5 [Mo] + 9.7 [Ni] + 7.9 [B] × 10 ３ +2.5] ×[C]; wherein [X] is the mass percent of the alloying element.

2. In mass percent, 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%, 0.0 0.001%≦Ca≦0.01%, 0.0001%≦Mg≦0.01%, 0.0001%≦Z At least one of r≦0.01% and 0.0001%≦REM≦0.01% is further satisfied. The steel sheet for hot stamping according to claim 1, further comprising:

3. 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+RE The steel sheet for hot stamping according to claim 2, wherein M≦0.30%.

4. The steel sheet substrate of the hot stamping steel sheet is, in mass percent, the balance being Fe and unreacted The steel sheet for hot stamping according to claim 1, which has free impurities.

5. The steel sheet substrate of the hot stamping steel sheet is, in mass percent, the balance being Fe and unreacted The steel sheet for hot stamping according to claim 3, which has free impurities.

6. 10. The method of claim 1, wherein Mf≧235° C. and / or h≦8.3≦h≦13.

5.

5. The steel sheet for hot stamping according to claim 4.

7. The steel sheet for hot stamping according to claim 6, wherein 8.4≦h≦12.

5.

8. The steel sheet for hot stamping according to claim 7, wherein 8.5≦h≦11.

5.

9. 0.28%≦C≦0.345%, and / or 0.15%≦Al≦0.38% The steel sheet for hot stamping according to claim 1 or 4,

10. The steel sheet for hot stamping according to claim 9, wherein 0.28%≦C≦0.31%.

11. The steel sheet for hot stamping according to claim 1 or 4, wherein Al / N≧65.

12. 0.80%≦Mn≦1.45%, and / or Si+Cr≦0.70% The steel sheet for hot stamping according to claim 1 or 4.

13. 0.10%≦Si≦0.40%, 0.01%≦Cr≦0.40%, and S The steel sheet for hot stamping according to claim 1 or 4, wherein i + Cr ≤ 0.50%.

14. 0.02%≦Nb+Ti≦0.06%, and / or 0.11%≦V≦0.2 5. The steel sheet for hot stamping according to claim 1, wherein the Cr content is 0%.

15. 2. The steel plate substrate has an average 10-point Vickers hardness of not more than 300 HV0.

3.

5. The steel sheet for hot stamping according to claim 4.

16. The microstructure of the steel sheet substrate has, in area percentage, martensite + bainite ≦ 30% The hot stamping material according to claim 1 or 4, wherein the remainder is ferrite and pearlite. Steel plate for construction.

17. 5 to 20 μm of aluminum alloy is coated on at least one surface of the steel plate substrate. The steel sheet for hot stamping according to claim 1 or 4,

18. The steel sheet is composed of a steel substrate and an outer layer from the inside to the outside, and the steel substrate is 0.27%≦C≦0.35%, 0.70%≦Mn≦1.55%, 0.10%≦Si ≦0.60%, 0.01%≦Cr≦0.70%, 0.001%≦B≦0.01%, 0.1 1%≦Al≦0.39%, N≦0.006%, 0.001%≦Nb+Ti≦0.1%, 0 .05%≦V≦0.20%, 0.001%≦P≦0.100%, 0.0001%≦S≦0 100%, Fe≧95%, and unavoidable impurities; Here, the martensitic transformation finish temperature Mf of the steel plate substrate satisfies Mf≧230°C. ;and The hardenability coefficient h of the steel plate substrate is calculated by the following formula and satisfies h≧8.0: h=[6.9[Mn] ２ +3.2[Si]+22.6[Cr] ２ +23.1[Mo]+( 13.0 [Cr] + 2.5 [Mo] + 9.7 [Ni] + 7.9 [B] × 10 ３ +2.5 ]×[C]; Wherein [X] is the mass percent of the alloying element. to.

19. In mass percent, 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%, 0.0 0.001%≦Ca≦0.01%, 0.0001%≦Mg≦0.01%, 0.0001%≦Z At least one of r≦0.01% and 0.0001%≦REM≦0.01% is further satisfied. The hot stamped molded component of claim 18, further comprising:

20. 0.0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+RE 20. The hot stamped component of claim 19, wherein M≦0.30%.

21. The steel sheet substrate of the hot stamping steel sheet is, in mass percent, the balance being Fe and unreacted 20. The hot stamped molded component of claim 18, having a free impurity content.

22. The steel sheet substrate of the hot stamping steel sheet is, in mass percent, the balance being Fe and unreacted 21. The hot stamped molded component of claim 20, having a free impurity content.

23. 17. Made from the steel sheet for hot stamping according to any one of claims 2 to 16.

22. The hot stamped component of claim 18 or 21.

24. 24. The hot stamped component of claim 23, wherein the hot stamp The outer layer of the punched component is composed of a decarburized layer of 1 to 20 μm, The C content in the hot stamped steel sheet does not exceed 50% of the C content in the center of the steel sheet substrate. Pre-assembled components.

25. Claim 18 or 21, which is made from the steel sheet for hot stamping according to claim 17.

2. The hot stamping component according to claim 1, wherein the hot stamping component The outer layer of the component is an interdiffusion layer of 4 to 15 μm, and Fe and Al outside the interdiffusion layer. and the Fe content of the interdiffusion layer is , a mass percentage of which is 70% or more.

26. The microstructure of the steel sheet substrate is less than 4% bainite, less than 3% austenite, and ferrite. The balance is dislocation martensite, and V, Nb or Ti.

18. The fine alloy carbides and / or nitrides formed by 22. A hot stamping molded component according to claim 21.

27. 27. The method of claim 26, wherein the sum of bainite + austenite + ferrite does not exceed 5%. Hot stamping molded components.

28. The microstructure of the steel plate substrate is formed by V, Nb, or Ti in area percent. The alloy contains 0.05 to 0.3% of fine alloy carbides and / or nitrides with an average particle size of 2 to 20 nm. The hot stamped molded component of claim 26 .

29. The hot stamping component has a yield strength of 1200 to 1450 MPa, Tensile strength is 1750-2100 MPa, elongation is 5% or more, and breaking strain is 0.22-(UTS The hot stamping method according to claim 18 or 21, wherein the hardness is 0.05-0.05 (-1700) / 5000 or more. Molded components.

30. The hot stamping component has a yield strength of 1300 to 1600 MPa, Tensile strength is 1700 to 2050 MPa, elongation is 5% or more, and breaking strain is 0.22 or more.

22. The hot stamped component according to claim 18 or 21.

31. The hot stamping component has an Al / N ratio of 65 or more and a yield strength of 1230 to 1420 MPa, tensile strength 1830-2030 MPa, elongation 5% or more, breaking strain 22. The composition according to claim 18, wherein the viscosity is 0.23-(UTS-1750) / 5000 or more. Hot stamping molding components.

32. A method for producing a steel sheet for hot stamping, comprising: The steel plate base material of the steel plate is, in mass percent, 0.27%≦C≦0.35%, 0.70%≦Mn ≦1.55%, 0.10%≦Si≦0.60%, 0.01%≦Cr≦0.70%, 0.0 0.01%≦B≦0.01%, 0.11%≦Al≦0.39%, N≦0.006%, 0.00 1%≦Nb+Ti≦0.1%, 0.05%≦V≦0.20%, 0.001%≦P≦0.1 00%, 0.0001%≦S≦0.100%, Fe≧95% and inevitable impurities, The martensitic transformation finish temperature Mf of the steel plate substrate satisfies Mf≧230°C, and The hardenability coefficient h of the steel plate substrate is calculated by the following formula, and satisfies h≧8.0: h=[ 6.9[Mn] ２ +3.2[Si]+22.6[Cr] ２ +23.1[Mo]+(13. 0 [Cr] + 2.5 [Mo] + 9.7 [Ni] + 7.9 [B] × 10 ３ +2.5]×[ C]; where [X] is the mass percent of the alloying element; The method comprises: a) Annealing treatment: Before coating, the steel sheet substrate is heated to a temperature of 740 to 870°C, and then annealed for 30 to 120 minutes. Incubate for 300 seconds, where the dew point is -30 to 5°C; b) Heat retention treatment: The annealed steel sheet substrate is cooled to a predetermined temperature in the range of 630 to 670°C. and keep warm for time t, （Ｆ＋Ｐ） ≦t≦100 seconds, where t （Ｆ＋Ｐ） teeth, The steel sheet substrate when heated to various annealing temperatures and then cooled to 630-670°C, % of ferrite and pearlite are produced; c) Hot-dip plating treatment: The steel sheet substrate kept warm in b) is immersed in a heated plating solution. A hot dip plating process is performed, wherein the composition of the plating solution is, in mass percent, Si: 9 to 1 2%, Fe: 2-3%, the balance being Al or Al alloy and unavoidable impurities. The temperature of the plating solution therein is maintained in the range of 630-670°C; d) After the steel sheet substrate is pulled out of the plating solution, and at least one of the steel sheet substrates is Before the plating solution on the other surface solidifies, remove the excess plating solution from at least one surface with an air filter. and removing the pre-plated layer on at least one of the surfaces by spraying with a spray gun. and e) The steel sheet substrate is cooled to room temperature and pre-plated with an aluminum alloy. obtaining a pre-plated steel sheet having a layer; A method comprising:

33. The steel sheet substrate of the hot stamping steel sheet has, in mass percent, 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%, and The following: 0.0001%≦Mg≦0.01%, 0.0001%≦Zr≦0.01%, and 0 0.0001%≦REM≦0.01%, wherein 0. 0001%≦W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≦0 33. The method for manufacturing a steel sheet for hot stamping according to claim 32, wherein the tensile strength is 0.30%.

34. The steel sheet substrate of the hot stamping steel sheet is, in mass percent, the balance being Fe and unreacted 34. The method for producing the steel sheet for hot stamping according to claim 32 or 33, which has an unavoidable impurity. How to do this.

35. t （Ｆ＋Ｐ） ≦t≦60 seconds, and / or t （Ｆ＋Ｐ） ≦35 seconds, claim Item 34. A method for producing a steel sheet for hot stamping according to item 32 or 33.

36. t （Ｆ＋Ｐ） ≦t≦40 seconds, and / or t （Ｆ＋Ｐ） ≦20 seconds, claim Item 34. A method for producing a steel sheet for hot stamping according to item 32 or 33.

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