WO2024104279A1 - 热冲压成形用钢板、热冲压成形构件及钢板制造方法 - Google Patents

热冲压成形用钢板、热冲压成形构件及钢板制造方法 Download PDF

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Publication number
WO2024104279A1
WO2024104279A1 PCT/CN2023/131146 CN2023131146W WO2024104279A1 WO 2024104279 A1 WO2024104279 A1 WO 2024104279A1 CN 2023131146 W CN2023131146 W CN 2023131146W WO 2024104279 A1 WO2024104279 A1 WO 2024104279A1
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Prior art keywords
hot stamping
steel sheet
steel plate
matrix
steel
Prior art date
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Ceased
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PCT/CN2023/131146
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English (en)
French (fr)
Chinese (zh)
Inventor
易红亮
孙正启
杨达朋
秦华杰
周澍
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Northeastern University China
Ironovation Materials Technology Co Ltd
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Northeastern University China
Ironovation Materials Technology Co Ltd
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Publication date
Application filed by Northeastern University China, Ironovation Materials Technology Co Ltd filed Critical Northeastern University China
Priority to EP23890713.3A priority Critical patent/EP4613900A1/en
Priority to JP2025527828A priority patent/JP2025541545A/ja
Priority to KR1020257019200A priority patent/KR20250109212A/ko
Publication of WO2024104279A1 publication Critical patent/WO2024104279A1/zh
Anticipated expiration legal-status Critical
Ceased legal-status Critical Current

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    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/12—Aluminium or alloys based thereon
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B32—LAYERED PRODUCTS
    • B32B—LAYERED PRODUCTS, i.e. PRODUCTS BUILT-UP OF STRATA OF FLAT OR NON-FLAT, e.g. CELLULAR OR HONEYCOMB, FORM
    • B32B15/00—Layered products comprising a layer of metal
    • B32B15/01—Layered products comprising a layer of metal all layers being exclusively metallic
    • B32B15/012—Layered products comprising a layer of metal all layers being exclusively metallic one layer being formed of an iron alloy or steel, another layer being formed of aluminium or an aluminium alloy
    • C—CHEMISTRY; METALLURGY
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/74—Methods of treatment in inert gas, controlled atmosphere, vacuum or pulverulent material
    • C21D1/76—Adjusting the composition of the atmosphere
    • C—CHEMISTRY; METALLURGY
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D3/00—Diffusion processes for extraction of non-metals; Furnaces therefor
    • C21D3/02—Extraction of non-metals
    • C21D3/04—Decarburising
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0257—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment with diffusion of elements, e.g. decarburising, nitriding
    • C—CHEMISTRY; METALLURGY
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
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    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C23C2/36—Elongated material
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite

Definitions

  • the invention relates to a steel plate for hot stamping, a hot stamping component and a method for manufacturing the steel plate.
  • Hot stamping components Lightweighting is an important way to achieve energy conservation and emission reduction in the vehicle industry.
  • the strength of hot stamping components obtained after hot deformation and cooling of hot stamping steel can reach more than 1500MPa.
  • hot stamping components When applied to the vehicle body, hot stamping components have a significant lightweight effect and ensure the collision safety of the vehicle. Therefore, the use of hot stamping components in vehicles has increased year by year.
  • CN106399837A provides a steel material or formed component with high toughness and delayed cracking resistance.
  • V nano-scale VC and/or (V, Ti, Nb) C composite carbides
  • the precipitation of these carbides refines the grains, thereby improving the yield strength and tensile strength of the material; on the other hand, it consumes carbon in austenite, thereby reducing the proportion of brittle twin martensite generated, improving the toughness of martensite; and on the other hand, it also improves the material's resistance to hydrogen-induced delayed cracking.
  • the patent adds more alloy elements that improve hardenability, such as Mn, Mo, etc., during alloy design.
  • the above-mentioned high alloy design is more likely to make the strip steel generate a hard and brittle martensitic structure in the processes of hot rolling, annealing and coating, causing cold rolling cracking or even strip breakage.
  • the strength of the material after annealing or coating is relatively high, resulting in problems such as difficulty in processing and difficulty in controlling the plate shape during the entire production process, which significantly increases the difficulty and cost of producing steel.
  • the patent also ignores the effect of coarse TiN particles on the toughness of the steel plate.
  • CN108374127A designed two Ti-free alloy compositions: 1. Low B content design, in which the role of B is replaced by a reasonable ratio of hardenability alloying 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, in which a certain amount of Al is added to replace the combination of Ti and N to avoid the precipitation of large-sized TiN particles, thereby To improve the toughness of hot stamping steel plates and their hot stamping components, that is, when 0.0005% ⁇ B ⁇ 0.005%, 0.4-1.0% Al is contained.
  • the hot stamping components obtained by the above alloy design can reach a yield strength of 1200-1800MPa, a tensile strength of 1500-2150MPa and an elongation of 7-10% after tempering, and a 99.5% confidence level of -40°C impact toughness ⁇ 45J ⁇ cm -2 .
  • CN108374127A also does not consider that the addition of high hardenability elements such as Mn, Cr, and Mo makes it easier to generate hard and brittle martensite during the production process, which is not conducive to cold rolling, annealing, coating and other processes and subsequent steel processing.
  • the addition of 0.4-1.0% Al increases the difficulty of continuous casting production and the problem of grain boundary oxidation on the surface of the steel plate, all of which will lead to an increase in the manufacturing cost of the steel plate.
  • CN106574348A specially controls the ratio of the content of four elements, namely C, Mn, Si and Cr, during alloy design, and ensures that a sufficiently low martensitic phase transformation start temperature Ms is obtained during the cooling period of the hot stamping process by adding 0.40-3% Mn, so that the tensile strength can be increased and the delayed cracking resistance can be improved.
  • the patent teaches that when Ni is concentrated on the surface of the plate or component in a specific form, 0.25-2% Ni can significantly reduce the sensitivity to delayed cracking.
  • Ni is an expensive alloying element, so the addition of Ni is not conducive to cost-effective hot stamping steel sheets and formed components thereof.
  • the patent increases the tensile strength by a low Ms temperature, but in actual production, the low Ms temperature will lead to the formation of hard and brittle twin martensite, which is not conducive to the improvement of the toughness of hot stamping components.
  • the present invention is made in view of the above problems existing in the prior art.
  • the steel sheet matrix of the hot stamping steel sheet of the present invention comprises, by mass percentage: 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 end temperature Mf of the steel plate matrix satisfies Mf ⁇ 230°C;
  • Mf is measured by experiment and the test method refers to "YB/T 5127-2018 Determination of critical point of steel" and is recorded in detail in the specific implementation method.
  • 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] 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).
  • the steel plate matrix for hot stamping further contains, in mass percentage, 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%.
  • the steel plate matrix of the hot stamping steel plate of the present invention contains, by mass percentage, the following: 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 balance is Fe and unavoidable impurities.
  • the steel sheet matrix of the hot stamping steel sheet comprises, by mass percentage, 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 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%, and 0.0001% ⁇ W+Mo+Ni+
  • 8.3 ⁇ h ⁇ 13.5 further preferably, 8.4 ⁇ h ⁇ 12.5, more preferably, 8.5 ⁇ h ⁇ 11.5.
  • Al/N ⁇ 65 Preferably, Al/N ⁇ 75.
  • the average value of the 10-point Vickers hardness of the steel plate substrate does not exceed 300 HV0.3.
  • the microstructure of the steel plate matrix comprises, in terms of area percentage, martensite+bainite ⁇ 30%, and the rest being ferrite+pearlite.
  • an aluminum alloy coating is coated on at least one surface of the steel plate substrate.
  • the hot stamping component of the present invention is composed of a steel plate matrix and an outer layer from the inside to the outside, and the steel plate matrix contains, by mass percentage: 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 Mf temperature of the steel plate substrate satisfies Mf ⁇ 230°C
  • the steel plate matrix of the hot stamping formed component further contains, in mass percentage, 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%.
  • the steel plate matrix of the hot stamping formed component contains, by mass percentage, the following: 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 balance is Fe and unavoidable impurities.
  • the hot stamping component is made of the aforementioned hot stamping steel sheet of the present invention.
  • the outer layer is composed of a decarburized layer of 1 to 20 ⁇ m, in which the hardness is less than that of the steel plate substrate. That is, along the thickness direction of the hot stamping formed component, 50% of the hardness at the center of the steel sheet matrix is used as the boundary between the decarburized layer and the steel sheet matrix from the outside to the inside.
  • the outer layer is a coating composed of an interdiffusion layer of 4 to 15 ⁇ m and an Fe and Al intermetallic compound layer outside the interdiffusion layer, wherein the Fe content of the interdiffusion layer is greater than or equal to 70% by mass. That is, along the thickness direction of the hot stamping formed component, from the outside to the inside, the Fe content is 70% as the boundary between the interdiffusion layer and the intermetallic compound layer.
  • the outer layer has a thickness of 5 to 40 ⁇ m.
  • the microstructure of the steel plate matrix of the hot stamping formed component is composed of, by area percentage, less than 4% bainite, less than 3% austenite, less than 3% ferrite, the remainder being dislocation martensite and less than 0.4% microalloy carbides and/or nitrides formed by V, Nb or Ti.
  • the sum of bainite+austenite+ferrite in the microstructure of the steel plate matrix of the hot stamping formed component does not exceed 5%.
  • the microstructure of the steel plate matrix of the hot stamping formed component contains 0.05-0.3% of microalloy carbides and/or nitrides formed by V, Nb, Ti, etc. in terms of area percentage, and the average size of the microalloy carbides and/or nitrides is 2-30 nm.
  • the yield strength of the hot stamping formed component is 1200-1450 MPa
  • the tensile strength is 1750-2100 MPa
  • the elongation is ⁇ 5%
  • the fracture strain is ⁇ 0.22-(UTS-1700)/5000.
  • the yield strength of the hot stamping formed component is 1300-1600MPa, the tensile strength is 1700-2050MPa, the elongation is ⁇ 5%, and the fracture strain is ⁇ 0.22.
  • the yield strength of the hot stamping formed component is 1230-1420MPa, the tensile strength is 1830-2030MPa, the elongation is ⁇ 5%, and the fracture strain is ⁇ 0.23-(UTS-1750)/5000.
  • the yield strength of the hot stamping formed component is 1350-1550 MPa
  • the tensile strength is 1750-2000 MPa
  • the elongation is ⁇ 5%
  • the fracture strain is ⁇ 0.23.
  • Another object of the present invention is a method for manufacturing a pre-plated steel sheet for hot stamping forming components, the method comprising:
  • Annealing treatment Before coating, heat the steel plate substrate to a temperature in the range of 740-870°C and keep it warm for 30-300s, with a dew point in the range of -30-5°C.
  • the steel plate matrix further comprises, in mass percentage, 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%.
  • the steel plate matrix contains Fe and inevitable impurities as the balance in terms of mass percentage.
  • the annealed steel plate substrate is cooled to a predetermined temperature in the range of 630-670°C and kept at the temperature for a time t, and t (F+P) ⁇ t ⁇ 100s, wherein t (F+P) is the time required for the steel plate substrate to generate 70% ferrite and pearlite when it is heated to different annealing temperatures and then cooled to 630-670°C;
  • c) hot dip treatment immersing the steel plate substrate in b) into a heated plating solution for hot dip plating, wherein the composition of the plating solution comprises, by mass percentage, 9-12% Si, 2-3% Fe and the remainder Al or Al alloy and inevitable impurities, and in this process, the plating solution temperature is maintained at a temperature within the range of 630-670° C.;
  • t (F+P) ⁇ t ⁇ 60 s; more preferably, t (F+P) ⁇ t ⁇ 40 s.
  • t (F+P) does not exceed 35 s, more preferably, t (F+P) does not exceed 20 s.
  • the plating solution temperature is not higher than the insulation temperature of the steel plate substrate in b).
  • the pre-plated steel sheet has an initial low carbon zone.
  • decarburization treatment dew point in the range of -30 to 5°C, preferably -28 to -5°C
  • the initial interface movement caused by the interdiffusion between the coating and the steel sheet substrate first occurs in the initial low carbon zone.
  • the presence of this initial low carbon zone causes very few carbon atoms in the newly generated interdiffusion layer to diffuse to the side of the base steel sheet and form enrichment, thereby reducing the formation of brittle martensite during subsequent cooling, which is beneficial to improving the toughness of hot stamping formed components.
  • the present invention makes the hot stamping forming components and hot stamping forming steel sheets of the present invention have Mf ⁇ 230°C and h ⁇ 8.0 by reasonably matching the alloy components, thereby obtaining products with high strength, high toughness and improved delayed cracking resistance.
  • Mf ⁇ 230°C avoids the formation of hard and brittle twin martensite during the phase transformation process, thereby improving the toughness and delayed cracking resistance of martensite.
  • h ⁇ 8.0 makes the hot stamping forming steel sheet have sufficient hardenability.
  • the properties of the hot stamping process can generate at least 95% martensite, preferably more than 97% martensite, to ensure high strength.
  • the combination of the two makes the hot stamping components have improved toughness and delayed cracking resistance while ensuring high strength.
  • the present invention aims to obtain a hot stamping component with a martensitic matrix that has high strength, good toughness and delayed cracking resistance.
  • martensite in steel can be divided into two types: dislocation type and twin type.
  • Dislocation martensite also known as low-carbon lath martensite, has a high formation temperature. Its high dislocation density causes carbon atoms to be mainly enriched on the dislocation line, forming dislocation gas masses, and does not form an oversaturated interstitial solid solution.
  • the martensite lattice lattice does not undergo obvious distortion, and is a body-centered cubic structure. Therefore, its high strength is mainly achieved through the interaction between carbon atoms and high-density dislocations.
  • twin martensite also known as high-carbon lamellar martensite, has a low formation temperature. A large number of carbon atoms are in the lattice in the oversaturated interstitial solid solution, resulting in obvious distortion of the lattice lattice, forming a body-centered square structure and a large number of fine twin substructures. Supersaturated carbon and fine twins make the twin martensite have ultra-high strength, but the fine twins hinder the slip of dislocations.
  • the present invention proposes to control the C content of the material and the content of other alloying elements in combination with the characteristics of the hot stamping production process, and adjust the martensitic phase transition temperature while ensuring hardenability to inhibit the formation of brittle twin martensite in the organization. Therefore, when designing the alloy ratio, it is not based on only high strength or only high toughness, but needs to consider both at the same time.
  • the present invention proposes to control the strength of austenite. This is because, during the martensitic phase transformation, the transformation of austenite to martensite is mainly a shear phase transformation, which occurs by the migration of martensite to the austenite interface, wherein the resistance to interface migration mainly comes from the strength of austenite.
  • the higher the austenite strength the greater the difficulty of phase interface migration, and the dislocation is difficult to slip, so that the interface migration has to be realized in the form of a twin substructure, thereby leading to the formation of twin martensite. Therefore, the present invention recognizes that in order to suppress the formation of twin martensite, the strength of austenite needs to be reduced. In the present invention, the strength of austenite is mainly reduced by the following two aspects.
  • the present invention controls the C content within the range of 0.27-0.35%.
  • the present invention realizes that the martensitic transformation temperature of steel directly affects the formation of twin martensite, and compared with the martensitic transformation start temperature, the martensitic transformation end temperature Mf has a more significant effect on the formation of twin martensite. This is because, although the martensitic transformation begins by generating dislocation martensite, as the temperature decreases, the generated martensite squeezes the original austenite, resulting in an increase in the strength of the austenite, making the dislocation martensite Slip is difficult to occur, so twin martensite begins to form.
  • the present invention proposes to control Mf.
  • high hardenability elements such as Mn and Cr are added to the steel plate in large quantities to improve the hardenability, so that the steel plate obtains more martensitic structure to have high strength.
  • the addition of the above-mentioned high content of alloying elements will reduce the Mf temperature, thereby making it easier to generate hard and brittle twin martensite during the phase transformation process, which reduces the toughness of the component and deteriorates the delayed cracking resistance; on the other hand, the excessive improvement of hardenability brought about by the high content of alloying elements will increase the difficulty of steel plate production and processing, which is not conducive to controlling production quality and cost.
  • the present invention optimizes the content of elements such as Mn and Cr and adds an appropriate amount of elements such as Al to increase the Mf temperature, and simultaneously obtains a higher Mf (Mf ⁇ 230°C) and a suitable hardenability coefficient h (h ⁇ 8.0), which can optimize the toughness of the martensite itself by obtaining as many dislocation martensite structures as possible while ensuring strength.
  • the present invention finds that under the synergistic effect of appropriate C content, hardenability and high Mf temperature, the strength of austenite during cooling is effectively controlled, so that when the martensite phase changes, it tends to generate as many strong and tough dislocation martensite with high dislocation density as possible, and inhibits the generation of brittle twin martensite, so that the final material has both high strength, high toughness and delayed cracking resistance.
  • the upper limit of the hardenability coefficient (h ⁇ 13.5) can be further limited to improve manufacturability and increase economic benefits.
  • h ⁇ 13.5 can improve the economy of material manufacturing, that is, it can avoid the generation of too much hard and brittle martensite phase in the production and manufacturing process (such as continuous casting, hot rolling, cold rolling and coating, etc.), thereby ensuring that the ingot, hot coil and coated finished coil have uniform and appropriate hardness, reducing the occurrence of problems such as slab cracks, cold rolling cracking or broken strips, and also improving the plate shape of the steel plate after cold rolling and coating, and the difficulty of subsequent processing (such as trimming, punching, leveling) of annealed or coated finished products is also reduced.
  • the proportion of martensite after hot stamping is not less than 95%, and it is ensured that the steel plate before hot stamping has good processing performance.
  • the hot rolled plate and coated/annealed plate in the steel plate production link have low tensile strength and low production cost.
  • adding an appropriate amount of Al element to the alloy and the ratio of Al/N are also another important consideration of the present invention.
  • the addition of a certain amount of Al can also form AlN with N, which has the effect of fixing N, thereby preventing the combination of N and B to ensure the role of B in improving hardenability.
  • AlN is formed at a relatively low temperature, the AlN particles formed are smaller in size and not easy to grow, and are not easy to cause serious damage to the toughness of the steel.
  • the addition of too much Al will significantly increase the austenite transformation end temperature of the steel, resulting in the inability to completely austenitize during hot stamping heating.
  • the present invention requires controlling the Al content Furthermore, the present invention finds that when the N content does not exceed 0.006%, controlling Al/N ⁇ 65 can ensure that the addition of Al can ensure the effect of B on improving hardenability while effectively avoiding the problem of increasing the production difficulty of the continuous casting link and the grain boundary oxidation problem on the surface of the steel plate.
  • the hot stamping formed component according to the present invention has sufficient toughness and delayed cracking resistance and high strength, showing good lightweight potential.
  • the hot stamping formed component has the following properties: yield strength 1200-1450MPa, tensile strength UTS: 1750-2100MPa, elongation ⁇ 5%, fracture strain ⁇ 0.22-(UTS-1700)/5000.
  • the higher fracture strain reflects that the component has sufficient toughness and delayed cracking resistance, which significantly reduces the risk of delayed cracking of the component during placement, processing, transportation, welding or loading after hot stamping.
  • FIG1 is a schematic diagram of the phase change expansion curve and Mf temperature calculation of the T1 pre-plated steel plate during the cooling process
  • FIG2 is a schematic diagram of the phase change expansion curve and t (F+P) calculation of the T1 pre-plated steel sheet during the insulation process when the annealing temperature is 820° C. and the insulation temperature is 647° C.;
  • FIG3 is a SEM microstructure of the steel plate substrate of the T1 pre-plated steel plate
  • FIG4 is a SEM microstructure of the steel plate substrate of the CT1 pre-plated steel plate
  • FIG5 is a SEM microstructure of the steel plate matrix of the untempered T1 hot stamping component
  • FIG6 is a TEM microstructure of the steel plate matrix of the untempered T1 hot stamping component
  • FIG. 7 is a diagram showing the relationship between the martensitic transformation end temperature Mf and the fracture strain of an untempered hot stamping component
  • FIG8 is a TEM microstructure of the steel plate matrix of the untempered CT1 hot stamping component
  • FIG9 is a SEM microstructure of the steel plate matrix of an untempered T5 hot stamping component.
  • FIG. 10 is a graph showing the relationship between the tensile strength and the fracture strain of an untempered hot stamped component.
  • the steel plate matrix of the first exemplary hot stamping formed component of the present invention comprises, by mass percentage: 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 hardenability coefficient of the steel plate matrix is h ⁇ 8.0 and Mf ⁇ 230°C.
  • the steel plate matrix of the second example hot stamping formed component of the present invention contains, in mass percentage, the following: 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, wherein the hardenability coefficient of the steel plate matrix is h ⁇ 8.0 and Mf ⁇ 230°C.
  • the steel plate matrix of the third example hot stamping formed component of the present invention contains, in mass percentage, the following: 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, wherein the hardenability coefficient of the steel plate matrix is h ⁇ 8.0 and Mf ⁇ 230°C.
  • the steel sheet matrix of the first, second and third exemplary hot stamping formed components contains Fe and inevitable impurities as the balance in terms of mass percentage.
  • the steel plate matrix of the first, second and third example hot stamping formed components also includes, by mass percentage, 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%.
  • the steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage, 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 at least one of the following: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.0 1% ⁇ 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%, 0.000
  • the steel sheet matrix of the hot stamping formed component of the present invention comprises, by mass percentage: 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 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%, and
  • the steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage, 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 at least one of the following: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.0 .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
  • the C element is the most effective alloying element for improving the strength of steel plates.
  • the C element as an important interstitial solid solution strengthening element, has a significant solid solution strengthening effect. Therefore, a C content lower than 0.27% cannot make the martensite have a sufficiently high strength. However, too much C is likely to lead to the formation of brittle twin martensite. Therefore, in order to ensure the toughness of martensite, the C content should not be too high. In addition, a high C content will also deteriorate the welding performance of the material. Finally, the C content in steel also significantly affects the phase change characteristics of steel, especially the Mf temperature.
  • the C content of the present invention is in the range of 0.27-0.35%.
  • the C content is preferably 0.28-0.345%, and the C content is more preferably 0.28-0.31%.
  • Mn has the effect of improving hardenability and significantly increasing the stability of austenite.
  • Mn can reduce the end temperature of austenite transformation, that is, expand the temperature range of complete austenitization.
  • the formation of ferrite and bainite can be delayed to promote martensitic transformation. Therefore, for hot stamping components, a certain amount of Mn is usually added to the alloy.
  • the Mn content of the present invention is in the range of 0.70 to 1.55%. Further, the Mn content is in the range of 0.80 to 1.45%.
  • N is an impurity element in steel, and especially for steel containing B, its combination with B will significantly reduce the effect of B in improving hardenability, so it is necessary to reduce the content of N as much as possible.
  • various processes for removing N all mean an increase in production costs, so, considering the production costs comprehensively, the N content in the present invention is controlled to be no more than 0.006%.
  • Al is an important element in the present invention.
  • Al is a strong deoxidizing element and is often used as a deoxidizer during steel smelting.
  • Al is a ferrite stabilizing element and can increase the Mf temperature, which is opposite to the effect of elements such as C, Mn, and Cr on the Mf temperature, thus helping to reduce the formation of hard and brittle twin martensite.
  • Al can form AlN with N, which has the effect of solidifying N, thereby preventing the combination of N and B and ensuring the effect of B in improving hardenability.
  • the formation temperature of AlN is relatively low, the AlN particles formed are relatively small in size and are not likely to cause serious damage to the toughness of the steel.
  • the Al content of the present invention is in the range of 0.11 to 0.39%, preferably in the range of 0.15 to 0.38%.
  • the present invention has found that the content range of the Al element can be further controlled by limiting Al/N ⁇ 65, while ensuring the role of B in improving hardenability while avoiding increasing the production difficulty of the continuous casting link and the problem of grain boundary oxidation on the surface of the steel plate.
  • Si is dissolved in the matrix to improve the strength of the matrix.
  • it can also be used as a deoxidizer in the steelmaking process. For this reason, more Si needs to be added.
  • more Si causes different degrees of grain boundary oxidation and decarburization to form on the surface of the steel plate at different positions during the cooling process of the steel coil after hot rolling, which affects the pickling effect of the steel coil and ultimately affects the surface quality of the finished product. Therefore, the Si content of the present invention is in the range of 0.10-0.40%.
  • Cr is an element that improves the hardenability of steel. At the same time, Cr has a significant effect on the anti-oxidation of steel and the prevention of surface decarburization. However, more Cr will make the oxide scale on the surface of the hot-rolled coil difficult to pickle, affecting the surface quality of the final product. Therefore, the Cr content of the present invention is in the range of 0.01 to 0.40%.
  • the sum of Si and Cr is not higher than 0.70%, and further, not higher than 0.50%.
  • B can be segregated at the austenite grain boundary, thereby inhibiting the formation of ferrite and significantly improving the hardenability of steel. Therefore, a certain amount of B is added to the component of the present invention. However, too high a B content will lead to boron embrittlement, which is not conducive to performance. Therefore, the B content of the present invention is within the range of 0.001 to 0.01%.
  • V, Nb, and Ti can form carbides in steel, which play a role in grain refinement and precipitation strengthening.
  • the formation of the above carbides consumes the carbon content in the matrix, which further improves the strength and toughness of the final product.
  • the excessive addition of these elements will not only not further improve the strength and toughness, but also lead to a significant increase in production costs.
  • Ti and N have a strong binding force and will form TiN with N in solid solution in steel. Although the combination of N and B elements is avoided and the role of B is guaranteed, since TiN precipitates at a higher temperature, it is easy to grow into coarse TiN particles during the subsequent cooling process of the slab, which is not conducive to the toughness of the steel plate.
  • the sum of the contents of Nb and Ti in the present invention is 0.001-0.10%, and the content of V is in the range of 0.05-0.20%. Further, the sum of the contents of Nb and Ti is in the range of 0.02-0.06%, and the content of V is in the range of 0.11-0.20%.
  • P is an unavoidable element.
  • P as a solid solution strengthening element can improve the strength of steel plates relatively cheaply.
  • the upper limit of the P content is not more than 0.100%, preferably not more than 0.050%.
  • the lower limit of the P content is not less than 0.001%, preferably not less than 0.004%.
  • the upper limit of the S content is not more than 0.100%, preferably not more than 0.015%.
  • the lower limit of the S content is not less than 0.0001%, preferably not less than 0.0005%, and more preferably not less than 0.001%.
  • the addition of W, Mo, Ni, Cu, and Co can improve the hardenability of steel, but the addition of these alloy elements will increase the cost of the alloy, so only an appropriate amount is added to the material.
  • Ni, Cu, and Co also have the benefit of improving the toughness of the material.
  • the content of each element is not less than 0.01%, the above-mentioned beneficial effect can be exhibited. Therefore, preferably, the lower limit of the content of the five elements is not less than 0.01%, respectively.
  • the upper limit of the content of the five elements is not more than 0.30%, respectively. In this case, while being able to improve toughness, it is ensured that the hardenability of the steel is less affected, and the machinability of the steel is guaranteed.
  • the content of Sn and Sb is not less than 0.005%, the wettability of the coating can be improved. Therefore, the lower limits of the contents of Sn and Sb are preferably not less than 0.005%, respectively. However, when more than 0.300% of Sn and/or more than 0.100% of Sb are contained, the toughness of the material will deteriorate. Therefore, the content of Sn is preferably not more than 0.300%, and the content of Sb is preferably not more than 0.100%.
  • REM Reare Earth Metal
  • Ca, Mg, and Zr can achieve the effect of refining inclusions by having a content of not less than 0.0001%, thereby improving the performance of the material. Therefore, the contents of Ca, Mg, Zr, and REM are preferably not less than 0.0001%, respectively.
  • the contents of Ca, Mg, Zr, and REM are preferably not more than 0.01%, respectively.
  • the sum of the contents of W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM elements is in the range of 0.0001% to 0.30%.
  • the steel plate matrix of the present invention may contain elements such as W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM, and the presence of these elements does not affect the solution of the technical problem of the present invention.
  • other components of the steel sheet matrix of the hot stamping component are not particularly limited.
  • elements such as As may be mixed from scrap, but if they are within a normal range, they will not affect the properties of the steel sheet matrix of the hot stamping steel sheet.
  • the hardenability coefficient h of the component of the present invention should be no less than 8.0 to ensure the generation of sufficient martensitic phase (no less than 95% by area percentage) and obtain a tensile strength of no less than 1700MPa.
  • the present invention proposes that the components also need to have a high Mf temperature. Specific explanation The explanation is as follows.
  • alloying elements such as C, Mn, and Cr will be increased.
  • These alloying elements are all stable austenitizing elements, which allow austenite to exist at a lower temperature, that is, to increase the strength of austenite. Since the martensitic phase transformation is a shear mechanism, when the strength of austenite is low, dislocation slip is easy to occur, and the martensitic phase transformation is mainly dislocation deformation, generating dislocation-type lath martensite; when the strength of austenite is high, the martensitic phase transformation is mainly twin deformation, generating twin martensite.
  • the present invention emphasizes that when developing hot stamping components with a strength of more than 1700MPa, under the condition of appropriate C content, other alloying elements need to be reasonably matched to simultaneously meet the requirements of hardenability and Mf temperature, so that high strength can be obtained while obtaining high toughness and improved delayed cracking resistance.
  • the content of elements such as Mn and Cr is adjusted to weaken their effect of reducing the end temperature of martensitic transformation on the basis of ensuring hardenability, and an appropriate amount of alloying elements such as Al that increase the Mf temperature is added to obtain a sufficiently high Mf temperature.
  • the present invention finds that with a C content of 0.27-0.35%, Mf ⁇ 230 ° C and hardenability of h ⁇ 8.0, the martensitic transformation will occur at a higher temperature, and the austenitic strength is effectively controlled, so that the martensitic transformation tends to generate as much dislocation martensite with high dislocation density as possible, and the generation of brittle twin martensite is suppressed, so that the final material has both high strength, high toughness and improved delayed cracking resistance. Furthermore, when the Mf temperature is above 235°C, the toughness and delayed cracking problems of the steel plate will be further improved.
  • the hardenability of the steel plate should not be too high, so as to avoid the generation of too much hard and brittle martensite phase in the production process of the steel plate (such as continuous casting, hot rolling and coating, etc.), ensure that the ingot, hot coil and coated finished coil have uniform and appropriate hardness, reduce the occurrence of problems such as slab cracks, cold rolling cracking or broken strips, and also improve the plate shape in the cold rolling and coating process. At the same time, the difficulty of subsequent processing (such as trimming, punching, and leveling) of the pre-coated steel plate is also reduced.
  • the hardenability coefficient h of the component of the present invention should be no more than 13.5, further, no more than 12.5, and further, no more than 11.5, so that the obtained hot stamping component has both high mechanical properties and improved manufacturability.
  • the hot stamping formed component is composed of a steel plate matrix and an outer layer from the inside to the outside.
  • a decarburized layer of a certain thickness will be formed on the surface during the hot stamping process, and the decarburized layer is the above-mentioned outer layer.
  • the presence of a decarburized layer of a certain thickness ensures that the component has good bending toughness.
  • a thicker decarburized layer not only damages the strength of the component, but also reduces the peak force of the VDA bending test.
  • the thickness of the decarburized layer of the present invention is 1 to 20 ⁇ m and the thickness of the decarburized layer is defined as follows: Through the Vickers hardness test, from the center of the steel plate matrix where the Vickers hardness is equal to 50% of the hardness to the surface of the hot stamping formed component. When there is an aluminum or aluminum alloy pre-plating layer on the surface of the steel plate used for hot stamping, during the hot stamping process, the elements in the aluminum or aluminum alloy pre-plating layer and the steel plate matrix will diffuse to form an interdiffusion layer and an Fe and Al intermetallic compound layer.
  • the outer layer of the hot stamping formed component is a coating composed of an interdiffusion layer and an Fe and Al intermetallic compound layer outside the interdiffusion layer and has a thickness of 5 to 40 ⁇ m, wherein the thickness of the interdiffusion layer is 4 to 15 ⁇ m and the interdiffusion layer is mainly ⁇ -Fe rich in Al and Si, wherein the Fe content is greater than or equal to 70% by weight.
  • the microstructure of the steel plate matrix of the component is composed of the following in terms of area percentage: less than 5% bainite, less than 3% austenite, less than 3% ferrite, and the rest is dislocation martensite and less than 0.4% carbides and/or nitrides formed by V, Nb, Ti, etc.
  • the less the non-martensitic phase in the microstructure, the better, and the sum of bainite, ferrite and residual austenite is required not to exceed 5%.
  • the average size of the microalloy carbides and/or nitrides formed by V, Nb, Ti, etc. needs to be controlled within the range of 2 to 30 nm, and the content is within the range of 0.05 to 0.4%.
  • the hot stamping formed components of the present invention exhibit ultra-high strength, good plasticity and toughness: a yield strength of 1200-1450MPa, a tensile strength UTS of 1750-2100MPa, an elongation of not less than 5%, and a fracture strain of not less than 0.22-(UTS-1700)/5000; preferably, a yield strength of 1230-1420MPa, a tensile strength UTS of 1830-2030MPa, and a fracture strain of not less than 0.23-(UTS-1750)/5000.
  • the hot stamping formed component of the present invention has the following properties: a yield strength of 1300-1600MPa, a tensile strength UTS of 1700-2050MPa, an elongation of not less than 5%, and a fracture strain of not less than 0.22; preferably, a yield strength of 1350-1550MPa, a tensile strength of 1750-2000MPa, and a fracture strain of not less than 0.23.
  • the tempering treatment is as follows: keep warm at a temperature of 150-230°C for 30-80min and then cool out of the furnace.
  • the present invention also provides a hot stamping steel plate for producing the above-mentioned hot stamping components, characterized in that the steel plate matrix of the hot stamping steel plate contains, in mass percentage: 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, wherein the Mf of the steel plate matrix satisfies Mf ⁇ 230°C and the hardenability coefficient h satisfies h ⁇ 8.0.
  • the steel sheet matrix of the hot stamping steel sheet further comprises, by mass percentage: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ At least one of 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%.
  • the steel sheet matrix of the hot stamping steel sheet contains Fe and inevitable impurities as the remainder in terms of mass percentage.
  • the hot stamping steel plate is a pre-plated steel plate, that is, the outer side of the steel plate substrate is coated with an aluminum or aluminum alloy coating.
  • the microstructure of the steel plate should be mainly ferrite and pearlite, and the formation of martensite phase should be avoided as much as possible, so that its proportion is less than 30%, preferably not more than 10%.
  • the formation of martensite during the processing of the steel plate can be avoided as much as possible by lowering the upper limit of h.
  • the structure of ferrite and pearlite makes the hardness of the steel plate lower, and the average value of the 10-point Vickers hardness does not exceed 300HV0.3, preferably, not more than 260HV0.3.
  • the thickness of the aluminum or aluminum alloy pre-plating layer is 5 to 20 ⁇ m.
  • the method for manufacturing the pre-plated steel sheet comprises:
  • the steel plate substrate having the above composition is heated to 740-870°C in an atmosphere of H2 and N2 with a H2 volume percentage of 2-12% and kept warm for 30-300s for annealing, with a dew point of -30-5°C; a dew point that is too high will cause serious oxidation of the steel plate surface, affecting the coating quality, while a low dew point will not bring about a beneficial surface decarburization effect;
  • t (F+P) is the time for the steel plate substrate with the above composition to be heated to different annealing temperatures and then rapidly cooled to 630-670° C. and kept at the temperature to generate 70% ferrite and pearlite;
  • the length of each stage of the coating and plating production line is fixed, and the residence time of the steel plate in each stage can be adjusted by adjusting the travel speed of the steel plate.
  • the insulation time t should not be less than the time t (F+P) for the steel plate matrix to be rapidly cooled to 630-670°C after being heated to different annealing temperatures to generate 70% ferrite and pearlite.
  • the insulation time t should not exceed 100s, preferably not more than 60s, and more preferably not more than 40s.
  • the composition of the plating solution comprises, by mass percentage, 9-12% Si, 2-3% Fe and the remainder Al or Al alloy and inevitable impurities, and in this process, the plating solution temperature is maintained at 630-670° C.; preferably, the plating solution temperature should not be higher than the insulation temperature in b) to prevent the plating solution temperature from decreasing after the steel plate is immersed in the plating solution, resulting in more Fe slag and affecting the quality of pre-coating;
  • the production method of the present invention improves production efficiency while reducing the occurrence of problems such as slab cracks, cold rolling cracking or strip breakage, and also improves the plate shape of the steel plate after cold rolling and coating, thereby reducing the difficulty of subsequent processing (such as trimming, punching, and leveling) of annealed or coated products.
  • the steel is prepared into experimental steel plates by the following process, namely:
  • Hot rolling The steel billet is heated to 1200°C and kept at this temperature for 2 hours, then hot rolled at 800°C to 1200°C, and coiled at 600°C to form a hot-rolled steel coil, which is then pickled to remove the oxide scale produced during the hot rolling process;
  • T1 to T8 are examples of the present invention, and CT1 to CT3 are comparative examples.
  • h is calculated according to the formula in the present invention, and Mf temperature is measured by using DIL805A phase change instrument with reference to the standard "YB/T 5127-2018 Critical Point Determination of Steel".
  • the specific test method is as follows:
  • T1-T8 and CT1-CT3 are subjected to the aforementioned method for manufacturing pre-plated steel sheets according to the parameters of Table 2, wherein the composition of the plating solution includes, by mass percentage: 9-12% Si, 2-3% Fe and the balance Al and inevitable impurities, and the plating solution temperature is maintained at 630-670°C. Afterwards, the steel plate substrate of the pre-plated steel sheet is subjected to Vickers hardness testing: the indenter load is 0.3kg ⁇ F, and the test result is the average value of 10 points.
  • the t (F+P) in Table 2 is obtained by simulating the corresponding method for manufacturing pre-plated steel sheets using a DIL805A phase change instrument, and FIG2 exemplarily illustrates the phase change expansion curve of the T1 pre-plated steel sheet during the insulation process when the annealing temperature is 820°C and the insulation temperature is 647°C and the calculation method of t (F+P) .
  • FIG3 and FIG4 show the microstructures of the steel plate substrates of the obtained T1 and CT1 pre-plated steel sheets.
  • the steel matrix of the T1 pre-plated steel plate is mainly composed of ferrite and pearlite, and the martensite content is less than 5%.
  • the typical microstructure of T1 is applicable to the steel matrix of T2 to T8 pre-plated steel plates. This microstructure corresponds to the hardness (227 to 275 HV0.3) of the steel matrix of T1 to T8 pre-plated steel plates.
  • the reason why the steel matrix is mainly composed of ferrite and pearlite is that the holding time before hot dip plating is longer than t (F+P) of the steel plate, so there is enough time to generate a large amount of ferrite and pearlite, reducing or avoiding the formation of martensite.
  • the steel sheet matrix of the CT1 pre-plated steel sheet is mainly composed of martensite and ferrite.
  • the content exceeds 30%.
  • Such a microstructure causes the hardness of the steel plate matrix to exceed 300HV0.3.
  • CT2 ⁇ CT3 pre-plated steel plates This is because the high alloy design of CT1 ⁇ CT3 makes the steel plate have a higher hardenability, among which the h value reaches 13.7 ⁇ 15.3, which makes it difficult to generate non-martensitic structures such as ferrite and pearlite during the annealing treatment of the coating process.
  • the t (F+P) time is relatively long, so that when the holding time is shorter than t (F+P) , the amount of ferrite and pearlite generated before hot dip plating is insufficient, and then in the cooling process after hot dip plating, the cooling rate is difficult to control, and more martensite will inevitably be generated, resulting in high hardness of the steel plate matrix, which is not conducive to subsequent processing.
  • CT3 it is noted that all its alloy element components are within the scope of the present invention, but the hardness of the steel plate matrix of the pre-plated steel plate is relatively high.
  • t (F+P) is preferably not greater than 35s, more preferably, t (F+P) is not greater than 20s and t is not greater than 60s.
  • the cold-rolled plates T1* and T4* i.e., without pre-plating
  • T1-T8 pre-plated steel plates and CT1-CT3 pre-plated steel plates were subjected to hot stamping flat plate simulation: heated to 920°C in a heating furnace for 300s, then the sample was transferred to the press for 8-12s, and after hot stamping, cooled to below 200°C at a cooling rate of 40°C/s to obtain the corresponding hot stamping formed components.
  • the above hot stamping formed components were tempered, and the tempering process was set to 170°C for 20min with reference to the actual baking process.
  • the T1-T8 and CT1-CT3 hot stamped components formed by hot stamping of pre-plated steel sheets they include a steel sheet substrate and an outer layer from the inside to the outside, and the outer layer is a coating composed of an interdiffusion layer and an Fe and Al intermetallic compound layer outside the interdiffusion layer.
  • the microstructure of the steel sheet substrate of each component was observed under a metallographic microscope and a scanning electron microscope, and the coating was subjected to EDS line scanning to obtain the change of the Fe content therein, so as to measure the thickness of the entire coating and the thickness of the interdiffusion layer, wherein the thickness of the interdiffusion layer was measured with a Fe content of 70% as the boundary between the interdiffusion layer and the Fe and Al intermetallic compound layer.
  • T1* and T4* hot stamped components formed by hot stamping of cold rolled sheets they include a steel sheet matrix and an outer layer from the inside to the outside.
  • the outer layer is a decarburized layer, and its thickness is defined as: from the surface of the hot stamped component to the position where the hardness is 50% of the hardness at the center of the steel sheet matrix.
  • the measurement method is as follows:
  • the tensile strength, elongation and maximum bending angle of the hot stamping formed structures without tempering and tempering were tested.
  • the final test results are the average of the three groups of test results.
  • the bending fracture strain test method is as follows: (1) Determine the VDA bending angle ⁇ peak of the component specimen by a static three-point bending test; (2) Based on the experimental results, select at least three groups of interruption bending angles ⁇ L (i.e., the bending angle of the component specimen under load) for interruption bending tests to ensure that ⁇ L ⁇ 50% ⁇ peak ; (3) Stop loading when the component specimen is bent to ⁇ L , and measure the bending angle ⁇ UL of the component specimen under unloading; (4) Place the unloaded component specimen under an optical microscope, and measure the inner and outer surface radii R i and R o of the most severe deformation zone; (5) Calculate the equivalent (plastic) strain ⁇ of the outer surface of the most severe deformation zone of the component specimen under unloading conditions of different ⁇ UL according to equation (1), that is, the equivalent strain of the outer surface of the most severe deformation zone of the component specimen when the component specimen is bent to ⁇ L , thereby establishing the ⁇ - ⁇ L relationship;
  • T-A refers to the thickness of the Al pre-plating layer
  • T-B refers to the thickness of the coating layer
  • T-D refers to the thickness of the interdiffusion layer
  • M refers to martensite
  • P refers to microalloy carbide.
  • the microstructures of all hot stamping parts are mainly martensite (accounting for no less than 95%), and the microalloy carbides are less than 0.3%.
  • the martensite content in the microstructures of T1-T4, T6-T8 and CT1-CT3 parts is The amount reaches more than 97%, which is almost a full martensite structure.
  • the nearly full martensite structure benefits from the sufficient hardenability of each component, ensuring high strength.
  • the typical microstructure of the T1 hot stamping formed 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 mainly composed of dislocation martensite with good toughness, and no twin martensite is generated.
  • the T1-T8, T1* and T4* components of the present invention have a yield strength of 1200-1400 MPa and a tensile strength of 1850-2030 MPa, a maximum bending angle of 45-53°, a fracture strain of 0.18-0.24, and have both high strength and high toughness.
  • the components of the present invention have both high strength and high toughness because the present invention obtains an Mf of not less than 230°C under the condition of satisfying the hardenability h ⁇ 8.0 through a reasonable alloy element ratio, such as the Mf of each embodiment T1 to T8 reaches 235 to 260°C. Therefore, after thermal deformation, at least 95% of the martensitic structure can be guaranteed during the cooling process of the mold, and no hard and brittle twin martensite is generated, so that the T1 to T8 components of the present invention show good toughness under high strength.
  • Figure 7 shows the relationship between the Mf temperature and the fracture strain of all the obtained untempered hot stamping formed components. It can be seen that with the increase of the Mf temperature, the toughness shows an increasing trend. Therefore, in order to ensure that the fracture strain is not less than 0.18, the Mf temperature of the present invention needs to be not less than 230°C, and the Mf temperature is preferably not less than 235°C.
  • the CT1 to CT3 components have a yield strength of 1200 to 1350 MPa and a tensile strength of 1900 to 2000 MPa, but the maximum bending angle does not reach 45° and the fracture strain is lower than 0.18, failing to have both high strength and high toughness.
  • the CT1-CT3 components have obtained sufficient high strength, but the excessive h value and low Mf temperature caused by their alloy design will make the CT1-CT3 components have more hard and brittle twin martensite, which is not conducive to the components to obtain sufficient toughness.
  • the Mf temperature of CT1 and CT2 is only 217°C and 225°C, because the Mn content of the two is as high as 1.79% and 1.60% respectively (exceeding the composition range of the present invention), so that even the high Al content of 0.59% and 0.45% (exceeding the composition range of the present invention) is not enough to compensate for the reduction of Mn on Mf.
  • CT3 composition is within the scope of the present invention, its element ratio is unreasonable, making h as high as 15.28 and Mf only 210°C.
  • the typical microstructure of the CT1 component is shown in Figure 8, which shows twin martensite, which corresponds to the expected effect of low Mf. Therefore, compared with the T1-T8 components of the present invention, although the CT1 component has high strength, the toughness performance is seriously insufficient.
  • the present invention requires that on the basis of each element satisfying the composition range of the present invention, each alloy element must be reasonably proportioned so that Mf ⁇ 230 ° C and h ⁇ 8.0. Further, it can be seen from Table 3 that h ⁇ 8.0 is sufficient to obtain a martensitic structure of not less than 95%. Considering manufacturability and economic benefits, the present invention preferably selects the hardenability coefficient to be 8.3 ⁇ h ⁇ 13.5, and more preferably, 8.5 ⁇ h ⁇ 11.5, such as T1-T8 components.
  • FIG. 9 shows a typical micrograph of an untempered T5 hot stamped component of the present invention.
  • the microstructure of the T1 component is a typical lath-shaped dislocation martensite structure, and the proportion of martensite is above 97%.
  • the Al/N of the T5 component is lower and fails to reach 65, resulting in N not fully combining with Al, but part of N combining with B, resulting in B's role in improving hardenability not being fully exerted, resulting in the formation of a small amount of ferrite, which is not conducive to ensuring high strength and good toughness.
  • the T1 component avoids the formation of ferrite by adding more Al (0.36%) to make Al/N ⁇ 65 to promote the combination of Al and N, thereby ensuring the role of B in improving hardenability in the steel plate.
  • the present invention preferably requires Al/N ⁇ 65 when 0.11% ⁇ Al ⁇ 0.39%.
  • FIG10 shows a relationship diagram between the fracture strain and the tensile strength of all the obtained untempered hot stamping formed components, wherein the fracture strain of the T1 to T8 and T1* and T4* components of the present invention is not less than 0.22-(UTS-1700)/5000, and further, if the Al/N ratio is not less than 65, the fracture strain of the embodiments of the present invention is not less than 0.23-(UTS-1750)/5000.
  • the final results of the tensile properties and VDA bending properties of the hot stamping formed components after tempering are shown in Table 4.
  • the T1-T8, T1* and T4* components of the present invention have a yield strength of 1320-1520 MPa and a tensile strength of 1780-1940 MPa, a maximum bending angle of 51-58°, and a fracture strain of 0.22-0.27, and have both high strength and high toughness.
  • the maximum bending angle of the CT1-CT3 components reached 47-50° and the fracture strain was 0.19-0.21 under the condition of a yield strength of 1350-1450 MPa and a tensile strength of 1800-1900 MPa.
  • the tempering treatment releases the internal stress of the CT1-CT3 components and improves their toughness to a certain extent, compared with the embodiments of the present invention, the CT1-CT3 components fail to have both high strength and high toughness.
  • the test method refers to the four-point bending loading test specified in ISO 7539-2:1995.
  • Five samples of hot stamping parts T1 ⁇ T8 and CT1 ⁇ CT3 that have not been tempered were taken for parallel experiments.
  • the solution was 0.1mol/N hydrochloric acid solution, and the load was applied according to 70 ⁇ 110% of the yield strength of the sample.
  • the maximum stress without cracking for 120 hours, that is, the ultimate stress and its ratio to the yield strength were recorded. The results are listed in Table 5.
  • the Mf of T1 ⁇ T3 and T6 ⁇ T8 is higher than 240°C, so that the ratio of ultimate stress to yield strength exceeds 60%, further improving the delayed cracking resistance of the component.
  • the Mf of T2 and T7 is higher than 255°C, so that the ratio of ultimate stress to yield strength exceeds 70%, further improving the delayed cracking resistance of the component.
  • the ultimate stress of components CT1 to CT3 is relatively low, and the ratio of ultimate stress to yield strength is less than 50%, and the delayed cracking performance is significantly worse than that of the embodiment of the present invention.
  • the hot stamping formed component of the present invention shows comprehensive superiority in overall performance, achieving high strength, high toughness and improved delayed cracking resistance.

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PCT/CN2023/131146 2022-11-14 2023-11-13 热冲压成形用钢板、热冲压成形构件及钢板制造方法 Ceased WO2024104279A1 (zh)

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KR1020257019200A KR20250109212A (ko) 2022-11-14 2023-11-13 열간 스탬핑 형성용 강철 시트, 열간 스탬핑 형성 부재 및 강철 시트 제조 방법

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Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN119571207A (zh) * 2025-02-07 2025-03-07 小米汽车科技有限公司 钢板基体、预镀层钢板及其制备方法、热冲压成型构件及其制备方法和汽车结构件

Families Citing this family (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN115478227B (zh) * 2022-11-14 2023-06-16 育材堂（苏州）材料科技有限公司 热冲压成形用钢板、热冲压成形构件及钢板制造方法
CN120249824A (zh) * 2025-06-04 2025-07-04 小米汽车科技有限公司 钢板基体、预镀层钢板及其制备方法、热冲压成型构件及其制备方法和汽车结构件

Citations (11)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20030173006A1 (en) * 2002-03-12 2003-09-18 Peppler William Joseph Steel product and method for manufacturing
CN104160046A (zh) * 2012-03-09 2014-11-19 株式会社神户制钢所 冲压成形品的制造方法和冲压成形品
CN104160051A (zh) * 2012-03-09 2014-11-19 株式会社神户制钢所 热压用钢板和冲压成形品以及冲压成形品的制造方法
CN104195455A (zh) * 2014-08-19 2014-12-10 中国科学院金属研究所 一种基于碳配分原理的热冲压烘烤韧化钢及其加工方法
CN106399837A (zh) 2016-07-08 2017-02-15 东北大学 热冲压成形用钢材、热冲压成形工艺及热冲压成形构件
CN106574348A (zh) 2014-07-30 2017-04-19 安赛乐米塔尔公司 用于模压淬火的钢板的制造方法和通过此方法获得的部件
CN107083475A (zh) * 2017-04-10 2017-08-22 钢铁研究总院淮安有限公司 一种热冲压成形防弹件的生产方法
JP2017179588A (ja) * 2016-03-29 2017-10-05 Ｊｆｅスチール株式会社 ホットプレス部材およびその製造方法
CN108374127A (zh) 2018-04-28 2018-08-07 育材堂（苏州）材料科技有限公司 热冲压成形用钢材、热冲压成形工艺及热冲压成形构件
CN113802065A (zh) * 2021-11-18 2021-12-17 育材堂（苏州）材料科技有限公司 热冲压成形构件、热冲压成形用钢板以及热冲压工艺
CN115478227A (zh) * 2022-11-14 2022-12-16 育材堂（苏州）材料科技有限公司 热冲压成形用钢板、热冲压成形构件及钢板制造方法

Family Cites Families (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP5756773B2 (ja) * 2012-03-09 2015-07-29 株式会社神戸製鋼所 熱間プレス用鋼板およびプレス成形品、並びにプレス成形品の製造方法
CN105543685A (zh) * 2015-12-25 2016-05-04 本钢板材股份有限公司 抑制热冲压过程氧化层脱落的热冲压成形用钢钢板及其生产方法
JP6583375B2 (ja) * 2016-09-28 2019-10-02 Ｊｆｅスチール株式会社 耐摩耗鋼板および耐摩耗鋼板の製造方法
CN110100030B (zh) * 2016-12-23 2021-04-20 Posco公司 弯曲加工性优异的超高强度热轧钢板及其制造方法
CN107829037B (zh) * 2017-09-15 2020-07-24 东北大学 热冲压成形用钢板、热冲压成形构件及梯度力学性能控制方法

Patent Citations (11)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20030173006A1 (en) * 2002-03-12 2003-09-18 Peppler William Joseph Steel product and method for manufacturing
CN104160046A (zh) * 2012-03-09 2014-11-19 株式会社神户制钢所 冲压成形品的制造方法和冲压成形品
CN104160051A (zh) * 2012-03-09 2014-11-19 株式会社神户制钢所 热压用钢板和冲压成形品以及冲压成形品的制造方法
CN106574348A (zh) 2014-07-30 2017-04-19 安赛乐米塔尔公司 用于模压淬火的钢板的制造方法和通过此方法获得的部件
CN104195455A (zh) * 2014-08-19 2014-12-10 中国科学院金属研究所 一种基于碳配分原理的热冲压烘烤韧化钢及其加工方法
JP2017179588A (ja) * 2016-03-29 2017-10-05 Ｊｆｅスチール株式会社 ホットプレス部材およびその製造方法
CN106399837A (zh) 2016-07-08 2017-02-15 东北大学 热冲压成形用钢材、热冲压成形工艺及热冲压成形构件
CN107083475A (zh) * 2017-04-10 2017-08-22 钢铁研究总院淮安有限公司 一种热冲压成形防弹件的生产方法
CN108374127A (zh) 2018-04-28 2018-08-07 育材堂（苏州）材料科技有限公司 热冲压成形用钢材、热冲压成形工艺及热冲压成形构件
CN113802065A (zh) * 2021-11-18 2021-12-17 育材堂（苏州）材料科技有限公司 热冲压成形构件、热冲压成形用钢板以及热冲压工艺
CN115478227A (zh) * 2022-11-14 2022-12-16 育材堂（苏州）材料科技有限公司 热冲压成形用钢板、热冲压成形构件及钢板制造方法

Non-Patent Citations (2)

* Cited by examiner, † Cited by third party
Title
See also references of EP4613900A1
ZHANG GUOQIANG ET AL.: "Progress in Research on Prediction Model of Hardenability of Steel", ACTA METALLURGICA SINICA, vol. 44, no. 4, 2019, pages 224 - 228

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN119571207A (zh) * 2025-02-07 2025-03-07 小米汽车科技有限公司 钢板基体、预镀层钢板及其制备方法、热冲压成型构件及其制备方法和汽车结构件
CN119571207B (zh) * 2025-02-07 2025-05-23 小米汽车科技有限公司 钢板基体、预镀层钢板及其制备方法、热冲压成型构件及其制备方法和汽车结构件

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