WO2022047956A1 - 带铝或铝合金预镀层的预镀层钢板、制造方法及热冲压成形构件 - Google Patents

带铝或铝合金预镀层的预镀层钢板、制造方法及热冲压成形构件 Download PDF

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WO2022047956A1
WO2022047956A1 PCT/CN2020/124190 CN2020124190W WO2022047956A1 WO 2022047956 A1 WO2022047956 A1 WO 2022047956A1 CN 2020124190 W CN2020124190 W CN 2020124190W WO 2022047956 A1 WO2022047956 A1 WO 2022047956A1
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Prior art keywords
steel sheet
hot
vda
base steel
range
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PCT/CN2020/124190
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English (en)
French (fr)
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易红亮
周澍
侯泽然
熊小川
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育材堂（苏州）材料科技有限公司
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Priority to KR1020237010917A priority Critical patent/KR20230061447A/ko
Priority to US18/024,384 priority patent/US20240026513A1/en
Publication of WO2022047956A1 publication Critical patent/WO2022047956A1/zh

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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
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D22/00—Shaping without cutting, by stamping, spinning, or deep-drawing
    • B21D22/02—Stamping using rigid devices or tools
    • B21D22/022—Stamping using rigid devices or tools by heating the blank or stamping associated with heat treatment
    • 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
    • C21—METALLURGY OF IRON
    • 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
    • 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
    • C21—METALLURGY OF IRON
    • 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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties 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 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
    • C21—METALLURGY OF IRON
    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
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    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
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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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    • 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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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
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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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    • C22C38/56—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.7% by weight of carbon
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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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    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
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    • 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
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C—CHEMISTRY; METALLURGY
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    • 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
    • C23C2/14—Removing excess of molten coatings; Controlling or regulating the coating thickness
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    • 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
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    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
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    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/14—Removing excess of molten coatings; Controlling or regulating the coating thickness
    • C23C2/16—Removing excess of molten coatings; Controlling or regulating the coating thickness using fluids under pressure, e.g. air knives
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    • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
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    • 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
    • C23C30/00—Coating with metallic material characterised only by the composition of the metallic material, i.e. not characterised by the coating process

Definitions

  • the present application relates to a pre-coated steel sheet with an aluminum or aluminum alloy pre-coated layer, a manufacturing method, and a hot stamped component.
  • a common test method for material toughness is the static three-point bending test (ie VDA bending test, VDA 238-100 standard), which evaluates the steel plate by detecting the bending angle (hereinafter referred to as the VDA bending angle) when the steel plate reaches the maximum bending load.
  • the toughness reflects the ability of the material to resist bending deformation failure.
  • a common test method to characterize the strength of materials is the room temperature tensile test (GB/T 228.1 standard), and the tensile strength reflects the ability of the material to resist tensile deformation failure.
  • EP2984198A1, CN102652177A and CN104769138A all make the final product have high tensile strength and good toughness by controlling the surface decarburization of the base steel plate.
  • EP2984198A1 relates to a plated thermoformed component.
  • a decarburization layer of 20-50 ⁇ m is formed on the surface of the base steel sheet at a dew point (eg, -15-5° C.) above -20° C. before coating, which helps to hinder the hot stamping process The tendency for microcracks to form in the base steel sheet.
  • a dew point eg, -15-5° C.
  • the carbon content is less than 0.01%, that is, complete decarburization
  • CN102652177A provides a method for manufacturing flat steel products with good formability.
  • CN102652177A shows that microstructural images of samples obtained using an annealing atmosphere dew point of -30°C show areas without decarburization.
  • this document teaches to control the dew point of the annealing atmosphere in the range of -20 to 60°C.
  • the decarburized edge layer is typically a ferrite structure with a maximum hardness of 75% of the central hardness of the flat steel product, avoiding the risk of cracks or nicks on the surface of the steel product during forming.
  • CN104769138A provides a method for making press hardened coated steel parts. Also, this patent finds that the formation of a decarburized zone with a depth p50% of 6-30 ⁇ m on the surface of the base steel sheet before the 22MnB5 pre-coating layer helps the final part to achieve high bendability, wherein the depth p50% is the carbon content equal to the base steel. The depth at 50% of the carbon content of the steel sheet. Furthermore, the data in this document show that the bend angle of the samples is undesirably less than 55°C at dew points below -15°C, and that the VDA bend angle decreases rapidly with decreasing dew point.
  • the document teaches that the dew point is not less than -15°C, that is, when the decarburization depth p50% is not less than 6 ⁇ m, 22MnB5 can obtain a critical bending angle higher than 55°.
  • the decarburized layer will significantly affect the ability of the hot stamped component to resist bending deformation failure, especially the maximum bending load (ie, the peak force corresponding to the VDA bending angle, hereinafter referred to as the VDA peak force), thereby affecting the component's crash safety. Therefore, it is unreasonable to only use the VDA bending angle and tensile strength to evaluate the crash safety of hot stamped components, and it is necessary to fully consider the influence of VDA peak force changes.
  • the present invention desires to obtain a pre-coated steel sheet with aluminum or an aluminum alloy pre-coated layer, a method for producing the same, and a hot stamping-formed member made therefrom.
  • the finally obtained hot stamped components not only have high toughness (VDA bending angle), but also have a high maximum bending load (VDA peak force), so that the Improve crash safety of hot stamped components.
  • the present invention provides a method of manufacturing a pre-coated steel sheet with an aluminum or aluminum alloy pre-coated layer, which enables a hot stamping formed member obtained from the pre-coated steel sheet to have excellent strength and toughness.
  • the coating method according to the present invention comprises:
  • the base steel plate is heated to the first temperature in the range of 740-880°C, preferably in the range of 740-820°C, in an ambient atmosphere of H2 and N2 with a volume percentage of 2-12 % H2 One temperature and hold for 30-300s, wherein, the carbon content C 0 of the base steel plate is in the range of 0.10-0.50%, and the manganese content is in the range of 0.50-10%, for the case of 0.10% ⁇ C 0 ⁇ 0.30% , the dew point of the controlled ambient atmosphere is in the range of -40 ⁇ -15°C, and for the case of 0.30% ⁇ C 0 ⁇ 0.50%, the dew point of the controlled ambient atmosphere is in the range of -36 ⁇ -12°C;
  • Hot dip plating the heated base steel sheet is cooled to a second temperature in the range of 610 to 680°C, and then immersed in a plating solution with a temperature of 610 to 680°C for hot dip plating;
  • the composition of the plating solution includes, by mass, 9 to 12% Si, 4% or less of Fe, the balance being Al, and unavoidable impurities.
  • the dew point of the ambient atmosphere is controlled to be in the range of -35 to -17°C, more preferably in the range of -31 to -19°C.
  • the dew point of the ambient atmosphere is controlled to be in the range of -30 to -15°C, more preferably in the range of -27 to -17°C.
  • the present invention provides a pre-coated steel plate coated with an aluminum or aluminum alloy pre-coated layer, the total thickness of the steel plate is 0.5-3.0 mm, preferably 0.7-2.3 mm, more preferably 0.8-2.0 mm, and the pre-coated steel plate includes A base steel sheet and a pre-coating of aluminium or an aluminium alloy on at least one surface of the base steel sheet,
  • the carbon content C 0 of the base steel sheet is in the range of 0.10-0.50%, and the manganese content is in the range of 0.50-10%;
  • the pre-plating layer thickness w 1 of the pre-plating layer is 5-20 ⁇ m, wherein the Al content is greater than or equal to 60% by mass;
  • An initial low carbon region exists in the base steel sheet adjacent to the interface between the base steel sheet and the pre-coating layer,
  • the base steel sheet contains the following components in mass percentage: 0.10-0.50% C, 0.50-10% Mn, 0-0.01% B, 0-0.4% Nb+Ti+V, 0.01-2% Si , 0.01-2% Al, 0.01-5% Cr+Ni+Mo+Cu and 0-2% Cr, 0-2% Ni, 0-2% Mo and 0-2% Cu, and The balance is Fe and inevitable impurity elements.
  • the present invention also provides a hot stamping forming member with aluminum or aluminum alloy coating, the thickness of the hot stamping forming member is 0.5-3.0mm, preferably 0.7-2.3mm, more preferably 0.8-2.0mm, from the inside to the outside
  • the hot stamping and forming components include:
  • the base steel sheet has a carbon content C 0 in the range of 0.10-0.50%, and a manganese content in the range of 0.50-10%;
  • An aluminum or aluminum alloy coating having a thickness of 10 to 26 ⁇ m and comprising: an interdiffusion layer adjacent to a base steel sheet, the interdiffusion layer having a thickness of 6 to 14 ⁇ m and comprising Al-containing ferrite, wherein the Fe content is greater than or equal to mass 70%; and an intermetallic compound layer of Fe and Al outside the interdiffusion layer;
  • the near-interface hardness HV 1 from the interface between the base steel sheet and the coating to within 6 ⁇ m of the base steel sheet is the core of the base steel sheet
  • the hardness is 0.65 to 1.07 times that of HV 2 and HV 2 is in the range of 400 to 550 HV.
  • HV 1 is 0.6 to HV 2 . 1.0 times and HV 2 is greater than 550HV.
  • the bending fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.283 and the VDA peak force is not less than the heat obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process 98% of the peak force of stamped-formed components; hot-stamped-formed components with a tensile strength higher than 1800MPa have a bending fracture strain of not less than 0.21 and a VDA peak force of not less than a non-decarburized pre-coating of the same composition and subjected to the same hot stamping process 97% of the peak force of the hot stamped part obtained from the steel sheet.
  • HV 1 is 0.70 to 1.0 times that of HV 2 ; for a hot stamped component with a tensile strength higher than 1800 MPa, HV 1 is 0.65 of HV 2 ⁇ 0.90 times.
  • the flexural fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.30 and the VDA peak force ratio is obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process.
  • the peak force of stamped-formed components is high; the bending fracture strain of hot-stamped-formed components with tensile strength higher than 1800MPa is not less than 0.23 and the VDA peak force is not less than that obtained from non-decarburized pre-coated steel sheets with the same composition and subjected to the same hot stamping process 99% of the peak force of the hot stamped part.
  • HV 1 is 0.75-0.95 times of HV 2 ; for hot-stamped components with tensile strength higher than 1800MPa, HV 1 is HV 2 0.68 to 0.85 times.
  • the bending fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.31 and the VDA peak force ratio is obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process.
  • the peak force of the stamped-formed member is at least 2% higher; the flexural rupture strain of the hot-stamped-formed member with a tensile strength higher than 1800 MPa is not less than 0.24 and the VDA peak force ratio is composed of the same composition and undergoes the same hot stamping process without decarburization pre-coating The peak force of the hot stamped part obtained from the steel sheet is high.
  • HV 1 and HV 2 are 10-point averages of Vickers hardness values measured using a load force of 5 g.
  • the present invention not only improves the toughness of the final formed member, but also avoids significantly reducing the tensile strength and maximum bending load of the member by controlling the initial low carbon region formed on the surface of the base steel sheet before pre-coating the steel sheet.
  • the control of the initial low carbon region is based not only on the expected VDA bending angle and tensile strength of the hot stamped part, but also on the VDA peak force, which will be explained in detail below.
  • the initial interface between the base steel sheet and the pre-plated layer will migrate to the side of the base steel sheet during the hot stamping process, so that compared with the pre-plated layer before hot stamping, the hot stamping
  • the thickness of the post coating is increased.
  • carbon atoms cannot diffuse to the side of the coating layer, so carbon atoms only diffuse to the side of the base steel sheet and diffuse to the side of the base steel sheet.
  • accumulation occurs in the base steel sheet, forming an obvious carbon-enriched area.
  • a brittle high-carbon martensite structure will be formed in the carbon-enriched area, and this layer of brittle high-carbon martensite will first crack during the static three-point bending test, which greatly impairs the toughness of the final formed component.
  • the strength of hot stamping materials is usually increased by increasing the carbon content in the alloy. As the carbon content in the base steel sheet increases, the damage to the toughness of the final component due to the high carbon martensite produced during hot stamping due to carbon enrichment will become more pronounced. In order to improve the toughness of high-strength aluminum-silicon-coated hot-stamped components, it is necessary to suppress or even eliminate carbon enrichment during hot-stamping.
  • the present invention proposes that in the production process of the pre-coated steel sheet, an initial low-carbon zone is formed on the surface of the base steel sheet before coating.
  • the initial interfacial movement caused by interdiffusion first occurs in the initial low carbon region.
  • the existence of the initial low-carbon region allows very few carbon atoms in the newly formed diffusion layer to diffuse to the side of the base steel sheet and form enrichment, thus greatly reducing the formation of brittle high-carbon martensite during subsequent cooling, thereby weakening the brittleness Damage of high carbon martensite to toughness of hot stamped parts.
  • the inventors have noticed that when materials or components with different thicknesses and strengths undergo bending deformation, their failures first occur on the outermost surface of the bending. This is because in this bending state, the outer surface is always affected by tensile stress. When the bending load reaches the limit, the outer surface will crack and cause fracture. At this time, the ultimate strain reached by the outer surface is called bending fracture strain, and the corresponding bending The angle becomes the VDA bend angle. Therefore, like the VDA bending angle, the bending fracture strain can also be used to characterize the toughness of a material or component, but the difference is that it is only related to the state of the outermost layer of the material or component, not the thickness of the material. Thus, the present application uses flexural fracture strain to characterize the toughness of the component, and the initial low carbon region is controlled by the desired flexural fracture strain.
  • the influence of the low carbon region on the tensile strength of the hot stamped part is often ignored in the prior art.
  • the alloyed coating on the surface is an intermetallic compound of Fe and Al, the hardness is as high as 800-1000HV, the brittleness is large, and the plasticity and toughness are poor.
  • the interdiffusion layer near the base steel plate is relatively soft high Al carbon-free ferrite, which has good plasticity and toughness, but low strength.
  • the coating on the surface of the component cannot play the role of bearing the tensile load, that is, the applied load is still carried by the base steel plate.
  • the stress state of the base steel plate is still consistent with that without decarburization.
  • the influence of the tensile strength of the component conforms to the classical mixing law, that is, the tensile strength of the hot stamped component will decrease linearly with the increase of the thickness of the low carbon region, but considering that the thickness of the low carbon region is generally several micrometers to tens of tens of The thickness of the base steel plate is several millimeters, and the thickness of the low-carbon region is much smaller than that of the base steel plate. Therefore, the existing technologies mostly ignore the influence of the thickness change of the low-carbon region on the tensile strength.
  • the present inventors found that it is unreasonable to use only the tensile strength and the VDA bending angle to evaluate the crash safety of the hot stamped member.
  • due to the difference in the microstructure and properties of the surface and the core of the hot stamping formed member made of aluminum-silicon coated steel plate due to the diameter of the bending indenter and the distance between the backup rolls. are relatively small, so when bending occurs, the component undergoes severe plastic deformation in the local micro area corresponding to the indenter and then fails. That is, the VDA bending test actually reflects the ability of the component to resist local plastic deformation failure.
  • the VDA bending test is also It is called the extreme sharp cold bending experiment. Considering various possible situations when vehicle collision occurs, it is unreasonable to evaluate the collision safety of hot stamping-formed components only by the tensile strength of the tensile test and the VDA bending angle of the VDA bending test. In order to evaluate the crash safety of a component, the inventors propose that a sufficiently high VDA peak force is also indispensable. However, the prior art has not paid attention to the effect of surface decarburization on the VDA peak force of hot stamped components.
  • the present invention finds that when the member undergoes VDA bending, the bending moment M 1 of the bending indenter satisfies the equation (1):
  • F is the bending load
  • L is the distance between the backup rollers
  • R is the diameter of the backup roller
  • r is the radius of the bending indenter
  • ⁇ is the bending angle
  • ⁇ (y) is the stress of the member when it is bent
  • W is the width of the member
  • t is the thickness of the member.
  • F peak is the VDA peak force when there is a low-carbon region on the surface of the base steel plate
  • ⁇ peak is the VDA bending angle when there is a low-carbon region on the surface of the base steel plate.
  • the VDA bending angle when there is a low carbon area on the surface of the base steel plate, the VDA bending angle will be improved, but when the VDA bending angle varies in a wide range, the VDA peak force will trend faster than the square relationship with the increase of the thickness of the low carbon area. decline.
  • the control of the decarburization zone should not only be based on the effect of surface decarburization on the tensile strength and VDA bending angle of the final component, but also based on the effect of surface decarburization on the VDA The effect of peak force. Therefore, the present invention proposes to control the initial low carbon region thickness on the surface of the base steel sheet (ie, control the dew point range) during the production of the pre-coated steel sheet, so that in the subsequent hot stamping process, due to the occurrence of the diffusion process, the pre-existing The initial low-carbon area is narrowed or even no longer exists. Therefore, the hot stamping formed component obtained according to the present invention has sufficient toughness, and the tensile strength and VDA peak force will not decrease significantly, thereby ensuring the collision safety of the component. sex.
  • the dew point can be taken from any range or any specific value within the range of -40 to -15°C, for example: -35 to -19°C, -31 to -20.1°C, -30 to -23°C, -29 ⁇ -20.1°C, -27 ⁇ -21°C, etc Any value such as °C, -26 °C, -26.4 °C, -27 °C, -28 °C, -32 °C, etc.
  • Figure 1 shows the micromorphology of the pre-coated steel sheet of T2 composition example A8;
  • Figure 2 shows the carbon distribution of the pre-coated steel sheets of T2 composition examples A8 and B2;
  • Figure 3 shows the relationship between the carbon content and the dew point of the initial low carbon region of the pre-coated steel sheets of the T1 and T2 compositions
  • Fig. 4 shows the micro-morphology and partial hardness indentation of the hot stamped-formed component obtained after hot stamping of the pre-coated steel sheet of T2 composition example A8;
  • Figure 5 shows the relationship between the decarburization degree (ie HV 1 /HV 2 ) and the dew point of the near-interface region of the base steel sheet of the hot stamped-formed component obtained after hot stamping of the pre-coated steel sheets of the T1 and T2 compositions;
  • Figure 6 shows the relationship between the bending fracture strain and the VDA bending angle of the hot stamping formed components obtained after the same decarburization treatment and hot stamping process for pre-coated steel sheets with different thicknesses of T1 composition
  • Figure 7 shows the relationship between the bending fracture strain and the VDA bending angle of the hot stamped-formed components obtained after hot stamping of pre-coated steel sheets with T1 and T2 compositions of 1.4 mm;
  • Fig. 8 shows the change of the bending fracture strain of the hot stamped-formed components obtained after hot stamping of pre-coated steel sheets with T1 and T2 compositions of 1.4 mm as a function of the ratio of HV 1 /HV 2 ;
  • Figure 9 shows the variation of the tensile strength of hot stamped-formed components obtained after hot stamping of pre-coated steel sheets with T1 and T2 compositions of 1.4 mm as a function of the ratio HV 1 /HV 2 ;
  • Figure 10 shows the variation of VDA peak force with the ratio of HV 1 /HV 2 for hot stamped-formed components obtained after hot stamping of pre-coated steel sheets with T1 and T2 compositions of 1.4 mm.
  • the improvement in toughness of the initial low carbon region is due to the fact that it counteracts or counteracts the carbon enrichment due to diffusion during hot stamping and the formation of a soft ferrite interdiffusion layer close to the base steel sheet. Therefore, compared to the existing pre-coated steel sheets with aluminum or aluminum alloy coatings, other conditions being equal, retaining a certain thickness of the initial low carbon area on the surface of the steel sheet will improve the hot stamping formed components made of the coated steel sheet toughness. It should be noted, however, that in order to avoid a significant decrease in VDA peak force due to decarburization by increasing toughness, the present inventors propose to control the thickness of the low carbon region, i.e.
  • the applicant also considered the influence of the carbon content of the base steel sheet on the existence of carbon enrichment in the subsequent hot stamping heating process.
  • the decarburization degree of the base steel sheet surface of the pre-coated steel sheet that is, the carbon content of the low-carbon region. Should not be too high.
  • the carbon enrichment phenomenon is relatively weak, so the degree of decarburization can be appropriately reduced, that is, the carbon content of the surface low-carbon area can be slightly higher.
  • the present invention provides a method of manufacturing a pre-coated steel sheet having an aluminum or aluminum alloy pre-coating layer in consideration of the relationship between the degree of decarburization and the dew point and the carbon content of the base steel sheet, which enables hot stamping obtained from the pre-coated steel sheet Formed components have excellent strength and toughness.
  • the coating method according to the present invention comprises:
  • the base steel plate is heated to the first temperature in the range of 740-880°C, preferably in the range of 740-820°C, in an ambient atmosphere of H2 and N2 with a volume percentage of 2-12 % H2 One temperature and hold for 30-300s, wherein, the carbon content C 0 of the base steel plate is in the range of 0.10-0.50%, and the manganese content is in the range of 0.50-10%, for the case of 0.10% ⁇ C 0 ⁇ 0.30% , the dew point of the controlled ambient atmosphere is in the range of -40 ⁇ -15°C, and for the case of 0.30% ⁇ C 0 ⁇ 0.50%, the dew point of the controlled ambient atmosphere is in the range of -36 ⁇ -12°C;
  • Hot dip plating the heated base steel sheet is cooled to a second temperature in the range of 610 to 680°C, and then immersed in a plating solution with a temperature of 610 to 680°C for hot dip plating;
  • the composition of the plating solution includes, by mass, 9 to 12% Si, 4% or less of Fe, the balance being Al, and unavoidable impurities.
  • the dew point of the ambient atmosphere is controlled to be in the range of -35 to -17° C. , more preferably in the range of -31 to -19°C.
  • the dew point of the ambient atmosphere is controlled to be in the range of -30 to -15°C, more preferably in the range of -27 to -17°C.
  • the present invention provides a pre-coated steel plate coated with an aluminum or aluminum alloy pre-coated layer, and the total thickness of the steel plate is 0.5-3.0 mm, preferably 0.7-2.3 mm, more preferably 0.8-2.0 mm, so
  • the pre-coated steel sheet includes a base steel sheet and a pre-coated layer of aluminum or aluminum alloy on at least one surface of the base steel sheet,
  • the carbon content C 0 of the base steel sheet is in the range of 0.10-0.50%, and the manganese content is in the range of 0.50-10%;
  • the pre-plating layer thickness w 1 of the pre-plating layer is 5-20 ⁇ m, wherein the Al content is greater than or equal to 60% by mass;
  • An initial low carbon region exists in the base steel sheet adjacent to the interface between the base steel sheet and the pre-coating layer,
  • the present invention also provides a hot-stamped component with aluminum or aluminum alloy coating, the thickness of the hot-stamped component is 0.5-3.0mm, preferably 0.7-2.3mm, more preferably 0.8-2.0mm, from the inside
  • the hot stamping and forming components include:
  • the base steel sheet has a carbon content C 0 in the range of 0.10-0.50%, and a manganese content in the range of 0.50-10%;
  • An aluminum or aluminum alloy coating having a thickness of 10 to 26 ⁇ m and comprising: an interdiffusion layer adjacent to a base steel sheet, the interdiffusion layer having a thickness of 6 to 14 ⁇ m and comprising Al-containing ferrite, wherein the Fe content is greater than or equal to mass 70%; and an intermetallic compound layer of Fe and Al outside the interdiffusion layer;
  • the near-interface hardness HV 1 from the interface between the base steel sheet and the coating to within 6 ⁇ m of the base steel sheet is the core of the base steel sheet
  • the hardness is 0.65 to 1.07 times that of HV 2 and HV 2 is in the range of 400 to 550 HV.
  • HV 1 is 0.6 to HV 2 . 1.0 times and HV 2 is greater than 550HV.
  • the bending fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.283 and the VDA peak force is not less than the heat obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process 98% of the peak force of stamped-formed components; hot-stamped-formed components with a tensile strength higher than 1800MPa have a bending fracture strain of not less than 0.21 and a VDA peak force of not less than a non-decarburized pre-coating of the same composition and subjected to the same hot stamping process 97% of the peak force of the hot stamped part obtained from the steel sheet.
  • HV 1 is 0.70 to 1.0 times that of HV 2 ; for a hot stamped component with a tensile strength higher than 1800 MPa, HV 1 is 0.65 of HV 2 ⁇ 0.90 times.
  • the flexural fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.30 and the VDA peak force ratio is obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process.
  • the peak force of stamped-formed components is high; the bending fracture strain of hot-stamped-formed components with tensile strength higher than 1800MPa is not less than 0.23 and the VDA peak force is not less than that obtained from non-decarburized pre-coated steel sheets with the same composition and subjected to the same hot stamping process 99% of the peak force of the hot stamped part.
  • HV 1 is 0.75-0.95 times of HV 2 ; for hot-stamped components with tensile strength higher than 1800MPa, HV 1 is HV 2 0.68 to 0.85 times.
  • the bending fracture strain of the hot stamped-formed member with a tensile strength in the range of 1300-1800 MPa is not less than 0.31 and the VDA peak force ratio is obtained from a non-decarburized pre-coated steel sheet having the same composition and undergoing the same hot stamping process.
  • the peak force of the stamped-formed member is at least 2% higher; the flexural rupture strain of the hot-stamped-formed member with a tensile strength higher than 1800 MPa is not less than 0.24 and the VDA peak force ratio is composed of the same composition and undergoes the same hot stamping process without decarburization pre-coating The peak force of the hot stamped part obtained from the steel sheet is high.
  • HV 1 and HV 2 are 10-point averages of Vickers hardness values measured using a load force of 5 g.
  • C is the most effective solid solution strengthening element in steel.
  • the carbon content needs to be greater than or equal to about 0.10%.
  • the carbon content exceeds 0.50%, its microstructure is mainly brittle high-carbon martensite, its ductility and toughness are poor, and its resistance to hydrogen embrittlement is significantly reduced. Therefore, the C content of the high-strength steel used in the present invention is between about 0.10-0.50%.
  • Mn about 0.50 to 10%
  • Mn is an element that increases the hardenability of steel to secure the strength of the steel.
  • the Mn content is less than about 0.50%, the hardenability of the steel material is insufficient, and it is difficult to obtain high strength.
  • the Mn content is too high, central band segregation is likely to occur in the steel, which adversely affects the ductility and toughness of the steel. Therefore, the upper limit of the Mn content of the high-strength steel used in the present invention is about 10%.
  • the improvement in the toughness of the obtained hot stamped components mainly comes from two aspects:
  • the interdiffusion layer in aluminum or aluminum alloy coating. Compared with the bare plate, the interdiffusion layer is a soft high Al ferrite, which has better toughness and plasticity, and the existence of this layer improves the toughness of the hot stamped component.
  • the thickness of the interdiffusion layer is optimal at 6-14 ⁇ m. When the thickness of the diffusion layer is less than 6 ⁇ m, the austenitization of the base steel sheet is insufficient, and the uniformity of the structure after cooling is poor, which affects the comprehensive performance of the final component.
  • the thickness of the diffusion layer is higher than 14 ⁇ m, it means that the heating time of hot stamping is too long or the temperature is too high, which will lead to the obvious growth of grains in the process of austenitization of the base steel plate, which will deteriorate the toughness of the final component;
  • the initial low carbon area of the pre-coated steel plate which reduces or offsets the carbon enrichment area formed by the accumulation in the base steel plate due to the occurrence of diffusion near the interface between the final coating and the base steel plate during the hot stamping process, thereby reducing or avoiding.
  • the carbon-enriched regions transform into brittle high-carbon martensitic structures during subsequent cooling, improving the toughness of the final hot stamped component.
  • the invention adopts the method of detecting the near-interface hardness of the base steel plate of the hot-stamped component and its ratio to the core hardness of the base steel plate to determine the decarburization or carbon enrichment degree of the near-interface of the base steel plate of the component after hot stamping.
  • the hardness ratio should not be too high. If it is too high, it means that there are still many brittle high-carbon martensite structures in the base steel plate, that is, the carbon enrichment phenomenon in the hot stamping process cannot be effectively reduced, which is not good for the toughness of the components. However, it should not be too low.
  • the hardness HV 1 has the following relationship with the core hardness HV 2 of the central area of the base steel plate:
  • the near-interface hardness HV 1 from the interface between the base steel sheet and the coating to within 6 ⁇ m of the base steel sheet is the core of the base steel sheet
  • the hardness is 0.65 to 1.07 times that of HV 2 and HV 2 is in the range of 400 to 550 HV.
  • HV 1 is 0.6 to HV 2 . 1.0 times and HV 2 is greater than 550HV.
  • HV 1 is 0.70 to 1.0 times that of HV 2 ; for a hot stamped component with a tensile strength higher than 1800 MPa, HV 1 is 0.65 of HV 2 ⁇ 0.90 times.
  • HV 1 is 0.75-0.95 times of HV 2 ; for hot-stamped components with tensile strength higher than 1800MPa, HV 1 is HV 2 0.68 to 0.85 times.
  • the base steel sheet with the composition shown in Table 1 was prepared, and the corresponding manufacturing process was as follows:
  • Hot rolling The billet is heated to 1200°C for 2 hours, then hot rolled at 800 ⁇ 1200°C, pickled to remove the oxide scale produced during hot rolling, and coiled below 700°C to form a hot rolled steel coil ;and
  • Cold rolling cold-rolling the pickled hot-rolled coil with a cold rolling reduction of 30-70% to obtain cold-rolled steel sheets with thicknesses of 1.2 mm, 1.4 mm and 1.8 mm.
  • the T1 component is 22MnB5, and the tensile strength of the hot stamped member using this alloy component is usually 1300 to 1800 MPa.
  • the T2 composition is 35MnB5V, and the tensile strength of the hot stamping component using this alloy composition can reach more than 1800MPa.
  • the description of the content of chemical elements in the text refers to the mass percentage.
  • the decarburization treatment of the above-mentioned cold-rolled steel sheet can be performed, for example, in the continuous annealing process of the cold-rolled steel sheet, and the ambient atmosphere is 5% H 2 +N 2 (volume fraction).
  • the process parameters in the decarburization process such as heating temperature, dew point and holding time, were adjusted, see Table 2.
  • the heated steel sheet is cooled to a range of 610-680° C. for hot-dip plating, and after that, the thickness of the pre-coating layer is controlled by blowing with an air knife, thereby obtaining a pre-coating steel sheet with pre-coating layers of aluminum or aluminum alloys of different thicknesses .
  • the composition of the plating solution includes by mass: 9-12% Si; and the balance is Al or Al alloy and inevitable impurities.
  • the pre-coating thickness of the pre-coating steel sheet was measured by scanning electron microscope (SEM), and the carbon content distribution from the interface between the base steel sheet and the pre-coating layer to about 18 ⁇ m in the base steel sheet was measured by GDOES glow discharge emission spectroscopy.
  • FIG. 1 shows the microstructure of the pre-coated steel sheet of Example A8, wherein the average pre-coated thickness is about 9.5 ⁇ m.
  • the distribution result of carbon element in the base steel sheet of the pre-coated steel sheet is shown in FIG. 2 .
  • Figure 2 shows the distribution of carbon elements in the base steel sheets of the example A8 and B2 pre-coated steel sheets.
  • the curve of Example B2 shows that the carbon content from the interface to 2 ⁇ m in the base steel sheet has reached about 90% of the carbon content C 0 of the base steel sheet.
  • the carbon content of Example A8 is only about 60% C 0 at 2 ⁇ m.
  • the carbon content of Example B2 was about 94.3% C 0 and the carbon content of Example A8 was about 69.1% C 0 .
  • the carbon content of Example B2 was about 96.8% C 0 and the carbon content of Example A8 was about 83.8% C 0 .
  • Example B2 The carbon content of Example B2 is higher than that of Example A8 because the dew point of Example A8 is higher than that of Example B2, so the degree of decarburization is increased, so that the initial low carbon region thickness of Example A8 is thicker than that of Example B2.
  • the dew point of Example B0 is the lowest, reaching -41°C, therefore, the degree of decarburization is the lightest, and the carbon content at 6 ⁇ m and 10 ⁇ m from the interface to the base steel sheet is about 97% C 0 and about 99% of the base steel sheet, respectively C 0 .
  • the dew point of Example B1 is the highest, reaching -5°C, so the degree of decarburization is the highest, and the carbon content at 6 ⁇ m and 10 ⁇ m from the interface to the base steel sheet is less than about 27% C 0 and about 59% C 0 of the base steel sheet, respectively.
  • the dew point is raised from -31°C to -17°C
  • the carbon content at 6 ⁇ m from the interface to the base steel sheet is in the range of 55.3 to 76.4% C 0
  • the carbon content at 10 ⁇ m is higher than High at 6 ⁇ m and in the range of 72.6 to 87.2% C 0
  • the dew point of Example B2 is -39°C
  • the surface of the base steel sheet is slightly decarburized
  • the carbon content from the interface to 6 ⁇ m and 10 ⁇ m in the base steel sheet is about 94% C 0 and about 97% C 0 , respectively.
  • Example B3 is -5°C
  • the surface decarburization of the base steel sheet is severe
  • the carbon content from the interface to the 6 ⁇ m and 10 ⁇ m in the base steel sheet is less than about 25% C 0 and about 55% C 0 of the base steel sheet, respectively.
  • FIG. 3 shows the relationship between the carbon content and the dew point of the initial low carbon region of the pre-coated steel sheets of the T1 and T2 compositions. It can be seen from the figure that as the dew point increases, the carbon content of the initial low-carbon region decreases, and accordingly, the thickness of the low-carbon region increases, that is, the degree of decarbonization increases. At the same time, compared with the T1 composition, due to the higher carbon content of the T2 composition, under the same dew point conditions, the decarburization effect of the T2 composition example is more obvious, so the carbon content of the initial low-carbon region of the T2 composition example is higher than that of the T1 composition. Low.
  • the hot stamped-formed member obtained from the pre-coated steel sheet includes a base steel sheet and a coating on its outer side.
  • the Vickers hardness of the near-interface region and the central region of the base steel plate was tested for the hot-stamped components, and the room temperature tensile properties and VDA bending properties were tested for the tempered components.
  • the Vickers hardness test method is as follows: select the indenter load F HV of 5g force, press into the surface of the sample from the interface between the base steel plate and the coating layer to within 6 ⁇ m of the base steel plate and the core of the base steel plate, and then remove the test force and the diagonal lengths d 1 and d 2 of the indentation were measured under SEM.
  • the corresponding near-interface hardness from the interface between the base steel sheet and the coating to within 6 ⁇ m of the base steel sheet and the core hardness of the base steel sheet are calculated using the following equation (8):
  • Figure 4 shows the local microstructure and partial hardness indentation of T2 composition example A8 after hot stamping, wherein the thickness of the coating is about 14.7-16.9 ⁇ m, and the thickness of the interdiffusion layer is about 6.7-7.5 ⁇ m.
  • test methods of room temperature tensile properties and VDA bending properties refer to GB/T 228.1 and VDA 238-100 standards respectively, and the sampling direction of the samples is the rolling direction.
  • three groups of samples were selected for testing of tensile properties and VDA bending properties, and the final results were the average of the three groups of test results.
  • the bending fracture strain testing method is as follows: (1) utilize the static three-point bending experiment to determine the VDA bending angle of the sample, ⁇ peak ; (2) based on the experimental results, choose at least three groups of interrupted bending angles ⁇ L (that is, the sample is in the load-bearing state) (3 ) Stop loading when the sample is bent to ⁇ L , and measure the bending angle ⁇ UL of the sample in the unloaded state; (4) Set the The unloaded sample is placed under an optical microscope, and the inner and outer surface radii Ri and R o of the most severely deformed area are measured; (5) According to equation (9 ) , the most severe deformation of the sample under the unloaded state of different ⁇ UL is calculated The equivalent (plastic) strain ⁇ of the outer surface of the area, that is, the equivalent strain of the outer surface of the most severely deformed area of the sample when the sample is bent to ⁇ L , so as to establish the ⁇ - ⁇ L relationship; (6) According to the fitting results, The flexural fracture strain of the
  • the present application uses the ratio of the hardness of the near-interface region to the core hardness (ie HV 1 /HV 2 ) of the base steel plate of the hot stamped component to reflect the degree of decarburization.
  • FIG. 5 shows the relationship between the decarburization degree (ie HV 1 /HV 2 ) and the dew point of the near-interface region of the base steel sheet of the hot stamped-formed component obtained after hot stamping of the pre-coated steel sheets of the T1 and T2 compositions. It can be seen from the figure that with the increase of dew point, the ratio of HV 1 /HV 2 shows a downward trend, which is consistent with the decrease of carbon content with the increase of dew point as shown in Figure 3. The ratio of HV 1 /HV 2 corresponds to the degree of decarburization of the pre-coated steel sheet.
  • Example A0 the near-interface hardness to core hardness ratio of Example A0 is 1.05, showing that there is still carbon enrichment, the VDA bending angle and bending fracture strain have increased to 62.8° and 0.294, respectively. This is because although the dew point of Example A0 is low, the surface of the base steel sheet is only slightly decarburized, but the slight decarburization can reduce the carbon enrichment phenomenon caused by interdiffusion during the hot stamping process, which makes the VDA bending properties relative to the existing ones. There are further improvements in technology.
  • the ratios of the near-interface hardness to the core hardness of the samples A1 to A4 with the same thickness of 1.4 mm are in the range of 0.81 to 0.95, which means that after hot stamping, there is no carbon enrichment near the interface between the base steel plate and the coating. .
  • Example A0 the toughness of Examples A1-A4 after hot stamping is further improved, the VDA bending angle can reach 64.2-66.4°, and the bending fracture strain is 0.309-0.315. It can be seen that a certain degree of decarburization exists on the surface of the matrix of the pre-coated steel sheet, which can effectively reduce the carbon enrichment caused by diffusion during the hot stamping process, thereby reducing the formation of brittle martensite and improving the toughness, which is manifested as the VDA bending angle. and increase in flexural fracture strain. In contrast, the ratio of near-interface hardness to core hardness of example B0 is 1.12, and the carbon enrichment phenomenon is obvious.
  • Example B2 had a near-interface hardness to core hardness ratio of 1.03, showing a slight carbon enrichment.
  • the VDA bending angle of Example B2 is only 45.4°, and the bending fracture strain is about 0.201.
  • the ratio of the near-interface hardness to the core hardness of Examples A7-A11 is in the range of 0.70-0.93, correspondingly, the VDA bending angle is increased to 50.3-57.7°, and the bending fracture strain is also increased to 0.224-0.272, which shows that the carbon enrichment phenomenon has a significant impact on the hardness of the core.
  • the effect of toughness of hot stamped components has been eliminated.
  • Figure 6 shows the relationship between the bending fracture strain and the VDA bending angle of the hot stamping-formed components obtained from pre-coated steel sheets with different thicknesses of T1 composition after the same decarburization treatment and hot stamping process.
  • the VDA bending angle decreases approximately linearly, but the bending fracture strain remains approximately constant. This is because, for a certain material composition, theoretically the same treatment process will lead to the same near-surface state of the obtained hot stamping-formed components, and the VDA bending experiment mainly leads to the fracture of the outermost surface (ie, the near-surface layer) of the component , therefore, the flexural fracture strain does not change with the thickness of the member. Therefore, it is reliable to use the bending fracture strain to evaluate the crash safety of materials with different thicknesses.
  • Figure 7 shows the flexural fracture strain versus VDA bending angle of hot stamped-formed components obtained from 1.4 mm pre-coated steel sheets of T1 and T2 compositions after hot stamping. Under the same thickness, the bending fracture strain ⁇ of the hot stamping-formed components of the two compositions has a linear relationship with the VDA bending angle ⁇ peak , which satisfies the following equation (10):
  • ⁇ is 0.2824 when ⁇ peak is 60°, preferably ⁇ is about 0.30 when ⁇ peak is 63.2°, more preferably ⁇ is about 0.30 when ⁇ peak is 65° about 0.31 hours.
  • ⁇ is about 0.21 when ⁇ peak is 47°, preferably ⁇ is about 0.23 when ⁇ peak is 50.5°, more preferably ⁇ is about 0.24 when ⁇ peak is 52.3° .
  • the hot stamped component of T2 composition with a strength of up to 1900MPa has lower VDA bending angle and lower bending fracture strain, that is, the increase of strength is accompanied by the decrease of component toughness.
  • Figure 8 shows the flexural fracture strain of hot stamped-formed components obtained after hot stamping from pre-coated steel sheets of 1.4 mm T1 and T2 compositions as a function of the HV 1 /HV 2 ratio.
  • the ratio of HV 1 /HV 2 decreases, i.e., the degree of decarburization in the near-interface region increases
  • the bending fracture strain of hot stamped-formed components shows an upward trend, but the trend gradually slows down. This is because the bending fracture strain is only related to the near-surface state of the base steel plate, and decarburization treatment can significantly improve the toughness and plasticity of the near-surface layer.
  • the steel plate will have good toughness, that is, high flexural fracture strain.
  • further increasing the degree of decarburization in the near-interface region contributes less to the improvement of the toughness of the near-surface layer, resulting in a limited increase in the flexural fracture strain.
  • Figures 9 and 10 show the variation of tensile strength and VDA peak force with the ratio HV 1 /HV 2 of hot stamped-formed components obtained from pre-coated steel sheets of 1.4 mm T1 and T2 compositions after hot stamping.
  • the tensile strength of the hot stamped part decreases slightly as the ratio of HV 1 /HV 2 decreases, ie, the degree of decarburization in the near-interface region increases.
  • the HV 1 /HV 2 of the sample B1 dew point -5°C
  • the highest degree of decarburization is based on the sample B0 (dew point is -41°C, it can be seen from the above data that it is close to no decarburization).
  • the ratio is less than 0.5, the tensile strength is reduced by about 61 MPa compared with Example B0, and the reduction rate is about 3.94%.
  • the tensile strengths of Examples A0-A4 with a lower degree of decarburization than that of Example B1 decreased within 1% relative to Example B0, and such a small decrease was negligible.
  • the value of HV 1 /HV 2 of Example B3 (dew point of -5° C.) with the highest degree of decarburization is less than 0.5 based on Example B2 (dew point of -39° C.), and its tensile strength is similar to Compared with Example B2, it is reduced by about 67MPa, and the reduction rate is about 3.48%.
  • the tensile strength of Examples A7-A11 with a lower degree of decarburization than that of Example B3 decreased within 2% relative to Example B0, and such a small decrease was negligible.
  • Example B0 HV 1 /HV 2 ratio of 1.12
  • the sample B0 has obvious carbon enrichment phenomenon, and the brittle high-carbon martensite structure formed by carbon enrichment can cause cracking under low load, so the early stage of the sample is failure occurred.
  • the sample B0 has eliminated some carbon enrichment relative to no decarburization at all, that is, the peak VDA force of the specimen should be lower in the case of no decarburization at all.
  • the carbon enrichment of Example A0 is partially offset, so that the carbon enrichment phenomenon is slight, and its VDA peak force is increased relative to that of Example B0.
  • the carbon enrichment in Example A1 is completely offset, so the VDA peak force is further elevated relative to Example A0.
  • Example A1 The other Examples A2-A4 were further decarburized relative to Example A1, however the VDA peak force decreased gradually relative to Example A1 because further increasing decarburization would reduce the peak force with carbon enrichment already fully offset. Therefore, when the degree of decarburization is so large that the reduction of VDA peak force caused by decarburization is greater than the increase of VDA peak force caused by eliminating carbon enrichment, it will appear that the VDA peak force is lower than that without decarburization.
  • the VDA peak force of Example B1 (with a dew point of -5° C.) was significantly reduced compared to Example B0, by about 500 N, a reduction of about 5.52%. This is not expected.
  • the VDA peak force of the less decarburized Examples A7-A11 HV 1 /HV 2 ratio in the range of 0.7 to 0.93) relative to Example B2 (dew point of -39°C) was relative to Example B0 is also raised.
  • the degree of decarburization is so large that the decrease in VDA peak force caused by decarburization is greater than the increase in VDA peak force caused by eliminating carbon enrichment, the VDA peak force will be lower than without decarburization.
  • the VDA peak force of Example B3 (with a dew point of -5°C) was significantly reduced compared to Example B2, by about 936 N, a decrease of about 7.74%.
  • the present application proposes to control the initial decarburization degree, so that it can not only reduce or eliminate the brittle high-carbon martensite structure, but also ensure that the VDA peak force reduction caused by decarburization is not excessively greater than the VDA peak value caused by eliminating carbon enrichment. The force is increased, thereby improving the crash safety and weight reduction of the hot stamped part.
  • VDA peak force In order to ensure crash safety, components need to have high VDA peak force while having high toughness.
  • the VDA peak force is not less than 9000 N (not less than 98% of the VDA peak force close to the example B0 without decarburization), then The ratio of HV 1 /HV 2 should not be less than 0.65.
  • the present application requires the dew point to be in the range of about -40 to -15° C. to meet the desired properties of the hot stamped component.
  • the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 1a at 6 ⁇ m from the interface to the base steel sheet satisfies 53% C 0 ⁇ C 1a ⁇ C 0 , and the carbon content C at 10 ⁇ m 1b satisfies 75% C 0 ⁇ C 1b ⁇ C 0 and C 1b >C 1a .
  • the present application specifies that for the T1 composition, the VDA peak force is not less than 9150N (about 0.84% higher than the VDA peak force of the example B0 near no decarburization), so the HV1/HV2 ratio should be not less than 0.70.
  • the VDA bending angle is not less than 63.2°, then the bending fracture strain should not be less than 0.30, and correspondingly, the ratio of HV 1 /HV 2 should not be greater than 1.00.
  • HV 1 /HV 2 is in the range of about 0.7 to 1.0, correspondingly, the dew point is about -35 to -17 ° C and the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 1a at 6 ⁇ m from the interface to the base steel sheet satisfies 59% C 0 ⁇ C 1a ⁇ 90% C 0 , and the carbon content C 1b at 10 ⁇ m 77.5% C 0 ⁇ C 1b ⁇ 95% C 0 and C 1b >C 1a is satisfied.
  • this application specifies that for the T1 composition, the VDA peak force is not less than 9300N (about 2.5 % higher than the VDA peak force of the example B0 near no decarburization), so the HV1/HV2 ratio should be not less than 0.75.
  • the VDA bending angle is not less than 65°
  • the bending fracture strain should not be less than 0.31, and correspondingly, the ratio of HV 1 /HV 2 should not be greater than 0.95.
  • HV 1 /HV 2 is in the range of about 0.75 to 0.95, correspondingly, the dew point is about -31 to -19 ° C and the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 1a at 6 ⁇ m from the interface to the base steel sheet satisfies 64% C 0 ⁇ C 1a ⁇ 82% C 0 , and the carbon content C 1b at 10 ⁇ m 80% C 0 ⁇ C 1b ⁇ 91.5% C 0 and C 1b >C 1a is satisfied.
  • the VDA peak force is not less than 11800 N (not less than 97% of the VDA peak force of Example B2 without decarburization), so HV The 1 /HV 2 ratio should not be less than 0.6.
  • the VDA bending angle is not less than 47°.
  • the bending fracture strain should be at least not less than 0.21, and correspondingly, the ratio of HV 1 /HV 2 should not be greater than 1. Therefore, HV 1 /HV 2 is in the range of about 0.6 to 1.0.
  • the present application requires a dew point in the range of about -36 to -12° C. to meet the desired properties of the hot stamped part.
  • the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 2a at 6 ⁇ m from the interface to the base steel sheet satisfies 42%C 0 ⁇ C 2a ⁇ 87% C 0 , and the carbon content at 10 ⁇ m
  • the content C 2b satisfies 65% C 0 ⁇ C 2b ⁇ 95% C 0 and C 2b >C 2a .
  • the VDA peak force is not less than 12000 N (not less than 99% of the VDA peak force close to the example B2 without decarburization) and the VDA bending angle is not less than 50.5°, correspondingly, the bending fracture strain should not be less than 0.23, and the HV 1 /HV 2 ratio is in the range of about 0.65 to 0.90.
  • the present application requires a dew point in the range of about -30 to -15°C to meet the desired properties of the hot stamped part.
  • the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 2a at 6 ⁇ m from the interface to the base steel sheet satisfies 50%C 0 ⁇ C 2a ⁇ 75% C 0 , and the carbon content C 2b at 10 ⁇ m satisfies 70% C 0 ⁇ C 2b ⁇ 86% C 0 and C 2b >C 2a .
  • the VDA peak force is not less than 12200N (about 0.93% higher than the VDA peak force of Example B2 near no decarburization) and the VDA bending angle is not less than 52.3°, correspondingly, the bending fracture strain should not be less than 52.3°.
  • the HV 1 /HV 2 ratio is in the range of about 0.68 to 0.85. In this case, the present application requires a dew point in the range of about -27 to -17°C to meet the desired properties of the hot stamped part.
  • the carbon content of the pre-coated steel sheet should satisfy: the carbon content C 2a at 6 ⁇ m from the interface to the base steel sheet satisfies 55%C 0 ⁇ C 2a ⁇ 70% C 0 , and the carbon content C 2b at 10 ⁇ m satisfies 75% C 0 ⁇ C 2b ⁇ 85% C 0 .
  • the present invention controls the carbon content (ie, the degree of decarburization) of the initial low-carbon region of the base steel sheet of the pre-coated steel sheet by setting an appropriate decarburization process, so that the hot stamping formed member obtained according to the present invention not only has improved Toughness, but also high tensile strength and VDA peak force, improving crash safety of hot stamped parts.

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Abstract

本申请涉及带铝或铝合金预镀层的预镀层钢板、制造方法及热冲压成形构件。本发明的制造具有铝或铝合金预镀层的预镀层钢板的方法涉及对于0.10％≤C 0≤0.30％的情况，控制混合气氛的露点在-40～-15℃范围内，并且对于0.30％＜C 0≤0.50％的情况，控制混合气氛的露点在-36～-12℃范围内。本发明还公开了由上述方法获得的涂覆有铝或铝合金预镀层的预镀层钢板及由该预镀层钢板获得的热冲压成形构件。该热冲压成形构件具有改善的韧性和VDA峰值力，提高了热冲压成形构件的碰撞安全性。

Description

带铝或铝合金预镀层的预镀层钢板、制造方法及热冲压成形构件 技术领域
本申请涉及带铝或铝合金预镀层的预镀层钢板、制造方法及热冲压成形构件。
背景技术
热冲压钢在汽车材料中的应用比例逐年递增，相应地，汽车行业对热冲压钢的强度和韧性的要求也越来越高。材料韧性的一种常用测试方法是静态三点弯曲试验(即VDA弯曲实验，VDA 238-100标准)，通过检测钢板达到最大弯曲载荷时的弯曲角(以下简称VDA弯曲角)的大小来评价钢板的韧性，反映材料抵抗弯曲变形失效的能力。同时，表征材料强度的一种常用测试方法是室温拉伸试验(GB/T 228.1标准)，抗拉强度反映了材料的抵抗拉伸变形失效的能力。
众所周知，随着材料强度的提升，其韧性会相应降低。因此，本领域技术人员已经在研究如何在保证热冲压钢具有高强度的同时改善其韧性。例如，EP2984198A1、CN102652177A及CN104769138A均通过控制基体钢板的表面脱碳使得最终产品具有高抗拉强度和良好的韧性。
EP2984198A1涉及一种带镀层的热成型构件。该文献教导在涂镀之前，在高于-20℃的露点(例如，-15～5℃)条件下，使基体钢板的表面形成20～50μm的脱碳层，其有助于阻碍热冲压过程中在基体钢板中形成微裂纹的趋势。同时，热冲压后基体钢板和其镀层的金属影响区之间仍然存在5～30μm的低碳区(碳含量低于0.01％，即完全脱碳)，其具有较好的延性，有助于消除在热成形和/或冷却过程中的应力，改善最终产品的塑性和韧性。
CN102652177A提供了一种制造具有良好成形性能的扁钢制品的方法。CN102652177A表明使用-30℃的退火气氛露点获得的试样的显微结构图显示出没有脱碳的区域。为了能够在扁钢表面获得一层10～200μm厚的可延展的脱碳边缘层，该文献教导控制退火气氛的露点在-20～60℃范围内。该脱碳边缘层典型为铁素体结构，最大硬度为扁钢制品的中心硬度的75％，避免了成形过程中在钢制品表面上发生裂缝或形成缺口的危险。
CN104769138A提供了一种用于制造经压制硬化的涂覆钢部件的方法。同样，该专利发现在22MnB5预涂镀层之前预先在基体钢板表面形成深度p50％为6～30μm的脱碳区有助于最终部件获得高的可弯曲性，其中，深度p50％为碳含量等于基体钢板碳含量的50％处的深度。此外，该文献的数据表明在露点低于-15℃时，样品的弯曲角非期望地小于55℃，且 VDA弯曲角随着露点的降低迅速下降。因此，为了保证期望的弯曲角，该文献教导露点不小于-15℃，即脱碳深度p50％不小于6μm时22MnB5才能得到高于55°的临界弯曲角。
以上现有技术均利用基体钢板的表面脱碳改善了最终热冲压成形构件的韧性，但注意到，这些现有技术并未意识到脱碳对热冲压成形构件在碰撞过程中抵抗变形失效能力的不利影响。这是因为在现有技术中，由于脱碳层厚度远低于基体钢板厚度，所以通常认为脱碳层对于抗拉强度的影响是可以忽略不计的。进而，通常也认为脱碳层对构件在碰撞过程中抵抗弯曲变形失效的能力的影响也是可以忽略不计的。然而，本发明人意外地发现，事实并非如此。相反，脱碳层会显著影响热冲压成形构件抵抗弯曲变形失效的能力，尤其是最大弯曲载荷(即，VDA弯曲角对应的峰值力，以下简称VDA峰值力)，进而影响构件的碰撞安全性。因此，仅利用VDA弯曲角和抗拉强度来评价热冲压成形构件的碰撞安全性是不合理的，需要充分考虑VDA峰值力变化带来的影响。
基于上述问题，本发明希望获得一种带铝或带铝合金预镀层的预镀层钢板、其制造方法、以及由其制成的热冲压成形构件。相对于现有技术存在的相近抗拉强度的热冲压成形构件，最终获得的热冲压成形构件不仅具有高的韧性(VDA弯曲角)，而且还具备高的最大弯曲载荷(VDA峰值力)，以改善热冲压成形构件的碰撞安全性。
发明内容
本发明提供了一种制造具有铝或铝合金预镀层的预镀层钢板的方法，其使得由预镀层钢板获得的热冲压成形构件具有优异的强度和韧性。根据本发明的涂镀方法包括：
a)脱碳处理：将基体钢板在H 2体积百分数为2～12％的H 2和N 2的环境气氛中加热至在740～880℃范围内，优选地在740～820℃范围内的第一温度并保温30～300s，其中，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内，对于0.10％≤C 0≤0.30％的情况，控制环境气氛的露点在-40～-15℃范围内，并且对于0.30％＜C 0≤0.50％的情况，控制环境气氛的露点在-36～-12℃范围内；
b)热浸镀：将加热后的基体钢板冷却到在610～680℃范围内的第二温度，之后浸入温度为610～680℃的镀液中进行热浸镀；
c)在基体钢板离开镀液后且在至少一个表面上的镀液凝固前，通过气刀吹扫来移除至少一个表面上多余的镀液以控制所述至少一个表面上的预镀层厚度w 1；及
d)将基体钢板冷却至室温以获得具有铝或铝合金预镀层的预镀层钢板，其中，预镀层厚度w 1为5～20μm，预镀层钢板的总厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm。
镀液组成包含以质量计：9～12％Si、4％以下的Fe、余量为Al以及不可避免的杂质。
优选地，对于0.10％≤C 0≤0.30％的情况，控制所述环境气氛的露点在-35～-17℃范围内，更优选地在-31～-19℃范围内。
优选地，对于0.30％＜C 0≤0.50％的情况，控制所述环境气氛的露点在-30～-15℃范围内，更优选地在-27～-17℃范围内。
本发明提供了一种涂覆有铝或铝合金预镀层的预镀层钢板，钢板总厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm，所述预镀层钢板包括基体钢板和在基体钢板的至少一个表面上的铝或铝合金的预镀层，
所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；
所述预镀层的预镀层厚度w 1为5～20μm，其中Al含量以质量计大于等于60％；
在所述基体钢板内邻接所述基体钢板与所述预镀层之间的界面存在初始低碳区，
(1)在0.10％≤C 0≤0.30％的情况下，
a)从所述界面至基体钢板内6μm处的碳含量C 1a满足53％C 0≤C 1a≤C 0；且
b)从所述界面至基体钢板内10μm处的碳含量C 1b满足75％C 0≤C 1b≤C 0且C 1b＞C 1a；
优选地，59％C 0≤C 1a≤90％C 0，同时，77.5％C 0≤C 1b≤95％C 0且C 1b＞C 1a；
更优选地，64％C 0≤C 1a≤82％C 0，同时，80％C 0＜C 1b≤91.5％C 0且C 1b＞C 1a；
(2)在0.30％＜C 0≤0.50％的情况下，
a)从所述界面至基体钢板内6μm处的碳含量C 2a满足42％C 0≤C 2a≤87％C 0；且
b)从所述界面至基体钢板内10μm处的碳含量C 2b满足65％C 0≤C 2b≤95％C 0且C 2b＞C 2a；
优选地，50％C 0≤C 2a≤75％C 0，同时，70％C 0≤C 2b≤86％C 0且C 2b＞C 2a；
更优选地，55％C 0≤C 2a≤70％C 0，同时，75％C 0＜C 2b≤85％C 0。
所述基体钢板以质量百分比计包含以下成分：0.10～0.50％的C，0.50～10％的Mn，0～0.01％的B，0～0.4％的Nb+Ti+V，0.01～2％的Si，0.01～2％的Al，0.01～5％的Cr+Ni+Mo+Cu且0～2％的Cr、0～2％的Ni、0～2％的Mo及0～2％的Cu，以及余量为Fe及不可避免的杂质元素。
本发明还提供了一种带铝或铝合金镀层的热冲压成形构件，热冲压成形构件的厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm，由内至外所述热冲压成形构件包括：
基体钢板，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；及
铝或铝合金镀层，其厚度为10～26μm并且包括：邻接基体钢板的相互扩散层，该相互扩散层的厚度为6～14μm并且包括含Al的铁素体，其中Fe含量以质量计大于等于70％；及在相互扩散层外侧的Fe和Al的金属间化合物层；
对于0.10％≤C 0≤0.30％且抗拉强度在1300～1800MPa范围内的热冲压成形构件，从基体钢板与镀层之间的界面至基体钢板6μm内的近界面硬度HV 1为基体钢板的心部硬度HV 2的0.65～1.07倍且HV 2在400～550HV的范围内，对于0.30％＜C 0≤0.50％且抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.6～1.0倍且HV 2大于550HV。优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.283并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的98％；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.21并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的97％。
优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.70～1.0倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.65～0.90倍。进一步优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.30并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.23并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的99％。
更优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.75～0.95倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.68～0.85倍。进一步优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.31并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高至少2％；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.24并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高。
HV 1和HV 2是使用5g的载荷力测得的维氏硬度值的10点平均值。
本发明通过控制在预涂覆钢板之前在基体钢板表面形成的初始低碳区，不但改善了最终成型构件的韧性，而且避免显著地降低构件的抗拉强度和最大弯曲载荷。对于初始低碳区的控制，不仅仅基于热冲压成形构件的预期VDA弯曲角和抗拉强度，而且还基于VDA 峰值力，这将详细解释如下。
在热冲压过程中，随着相互扩散层的形成，基体钢板与预镀层之间的初始界面在热冲压过程中会向基体钢板一侧迁移，使得相较于热冲压前的预镀层，热冲压后镀层厚度增加。在界面迁移的过程中，由于碳原子在含Al的铁素体和/或FeAl化合物中溶解度极低，导致碳原子不能向镀层一侧扩散，所以碳原子仅向基体钢板一侧发生扩散并在最终镀层/基体钢板界面附近在基体钢板内发生堆积，形成明显的碳富集区。在随后的冷却中，碳富集区会生成脆性高碳马氏体组织，这层脆性高碳马氏体组织在静态三点弯曲试验时会首先发生开裂，极大地损害了最终成型构件的韧性。此外，为满足轻量化的需求，通常通过增加合金中的碳含量来提高热冲压材料的强度。随着基体钢板中碳含量的提高，热冲压过程中由于碳富集问题生产的高碳马氏体对最终构件韧性的损害将更加明显。为改善高强度的铝硅镀层热冲压成形构件的韧性，需要抑制甚至消除热冲压过程中的碳富集现象。本发明提出在预涂镀钢板生产过程中先在其基体钢板表面形成初始低碳区再进行涂镀。在这种情况下，在后续热冲压过程中，相互扩散导致的初始界面移动首先发生在初始低碳区中。该初始低碳区的存在使得新生成的扩散层中极少的碳原子向基体钢板一侧扩散并形成富集，故而大幅减少了随后冷却时脆性高碳马氏体的生成，从而削弱了脆性高碳马氏体对热冲压成形构件韧性的损害。
此外，本发明人注意到，不同厚度和强度的材料或构件在发生弯曲变形时，其失效均首先发生在弯曲最外侧的表面上。这是因为在该弯曲状态下，外侧表面始终受拉应力影响，当弯曲载荷达到极限时，外侧表面发生开裂从而造成断裂，此时，外侧表面达到的极限应变称为弯曲断裂应变，对应的弯曲角度成为VDA弯曲角。因此，与VDA弯曲角相同，弯曲断裂应变也可以用来表征材料或构件的韧性，但不同的是，其仅与材料或构件最表层的状态有关，与材料厚度无关。故而，本申请使用弯曲断裂应变来表征构件的韧性，并且通过期望弯曲断裂应变来控制初始低碳区。
如前所述，现有技术中常常忽略低碳区对于热冲压成形构件的抗拉强度的影响。对于热冲压成形构件，由于表面的合金化镀层为Fe和Al的金属间化合物，其硬度高达800～1000HV，脆性大，塑性和韧性差，在热冲压过程中会发生开裂形成大量微裂纹。相对而言，靠近基体钢板的相互扩散层为相对较软的高Al无碳铁素体，塑性和韧性均很好，但强度低。因此，通常情况下，在拉伸实验中，构件表面的镀层无法起到承担拉伸载荷的作用，即所施加的载荷仍是由基体钢板进行承载。当基体钢板与镀层之间的界面至基体钢板内的近界面区内存在低碳区时，在拉伸过程中，基体钢板受力状态仍与无脱碳情况保持一致，因此，低碳区对构件的抗拉强度的影响符合经典的混合法则，即，热冲压成形构件的抗拉强 度将随着低碳区厚度的增加呈线性下降，但考虑到低碳区厚度一般为几微米到几十微米，而基体钢板厚度为几毫米，低碳区厚度远小于基体钢板厚度，因此现有技术大多忽略了低碳区厚度变化对于抗拉强度的影响。
然而，本发明人发现仅用抗拉强度和VDA弯曲角来评价热冲压成形构件的碰撞安全性是不合理的。一方面由于由铝硅镀层钢板制成的热冲压成形构件的表面与心部组织性能有差异，另一方面，在进行热冲压成形构件的VDA弯曲实验时，由于弯曲压头直径和支撑辊间距均较小，所以发生弯曲时，构件在压头对应的局部微小区域发生严重的塑性变形进而形成失效，即，VDA弯曲实验实际反映了构件抵抗局部塑性变形失效的能力，因此，VDA弯曲实验又被称为极限尖冷弯实验。考虑到车辆碰撞发生时的各种可能情况，仅用拉伸实验的抗拉强度和VDA弯曲实验的VDA弯曲角来评价热冲压成形构件的碰撞安全性是不合理的。为评价构件的碰撞安全性，本发明人提出足够高的VDA峰值力也是不可或缺的。而现有技术中还没有关注到表面脱碳对于热冲压成形构件的VDA峰值力的影响。
本发明发现，在构件发生VDA弯曲时，弯曲压头的弯矩M 1满足等式(1)：
Figure PCTCN2020124190-appb-000001
其中，F为弯曲载荷，L为支撑辊间距，R为支撑辊直径，r为弯曲压头半径，α为弯曲角。
同时，构件发生VDA弯曲时，变形区域的弯矩M 2满足等式(2)：
Figure PCTCN2020124190-appb-000002
其中，σ(y)为构件弯曲时受到的应力，W为构件宽度，t为构件厚度。
由于等式(1)和等式(2)的弯矩相等，当弯曲角α达到最大弯曲角α max(即基体钢板表面无脱碳时的VDA弯曲角)时，最大弯曲载荷F max(即VDA峰值力)满足等式(3)：
Figure PCTCN2020124190-appb-000003
对于一定板厚的热冲压成形构件，当基体钢板表面存在厚度为t 1的低碳区时，由于低碳区具有低强度、较好的塑性和韧性，所以在发生弯曲时，处于构件外表面的低碳区不会首先发生 失效，而是次表层达到材料的抗拉强度形成开裂。因此，由等式(3)可知，表面存在低碳区时的VDA峰值力F peak应满足等式(4)和(5)：
Figure PCTCN2020124190-appb-000004
Figure PCTCN2020124190-appb-000005
其中，F peak为基体钢板表面存在低碳区时的VDA峰值力，α peak为基体钢板表面存在低碳区时的VDA弯曲角。以t为1.4mm为例，由VDA238-100标准可知，L为3.3，R为15，r为0.4。而对于本领域技术人员，1500MPa的22MnB5的无镀层热冲压成型钢在基体钢板表面无脱碳时，其VDA弯曲角α max可达到60°。因此，等式(5)可转化为等式(6)：
Figure PCTCN2020124190-appb-000006
由等式(6)可知，当60°≤α peak≤105°时，f(α peak)始终不大于1，VDA峰值力F peak将满足等式(7)：
Figure PCTCN2020124190-appb-000007
即当基体钢板表面存在低碳区时，VDA弯曲角将得到改善，但VDA弯曲角在很大范围内变化时，VDA峰值力将随着低碳区厚度的增加以比平方关系更快地趋势下降。
综上，对于带铝或铝合金预镀层的热冲压成形构件，虽然低碳区的存在抑制甚至消除了热冲压过程中的碳富集现象，进而改善了构件的VDA弯曲角和弯曲断裂韧性，但随着低碳区厚度的进一步增加，热冲压成形构件的VDA峰值力将显著下降。
鉴于上述，为了改善热冲压成形构件的碰撞安全性，对于脱碳区的控制不仅仅要基于表面脱碳对最终构件的抗拉强度、VDA弯曲角的影响，而且还要基于表面脱碳对VDA峰值力的影响。因此，本发明提出在生产预镀层钢板时，通过控制基体钢板表面的初始低碳区厚度(即，控制露点范围)，使得在随后的热冲压工艺过程中，由于扩散过程的发生，预先存在的初始低碳区缩窄，甚至不再存在，因此，根据本发明获得的热冲压成形构件在具有充足韧性的同时，抗拉强度和VDA峰值力不会出现明显下降，从而保证了构件的碰撞安全性。
此外，本领域技术人员将理解，上述各个区间内的任一范围或任一值都适用于本发明。比如，露点可以取自在-40～-15℃范围内的任一范围或任一具体数值，例如：-35～-19℃、-31～-20.1℃、-30～-23℃、-29～-20.1℃、-27～-21℃等中的任一范围，或诸如-20.2℃、-21.5℃、-22.2℃、-22.8℃、-23.6℃、-24℃、-24.7℃、-25.2℃、-26℃、-26.4℃、-27℃、-28℃、-32℃等任一数值。
附图说明
本发明的实施例、特征和优点将从结合附图的随后描述变得清楚，在附图中：
图1示出了T2成分示例A8的预镀层钢板的微观形貌；
图2示出了T2成分示例A8和B2的预镀层钢板的碳元素分布情况；
图3示出了T1和T2成分的预镀层钢板的初始低碳区的碳含量与露点之间的关系；
图4示出了T2成分示例A8的预镀层钢板在热冲压后获得的热冲压成形构件的微观形貌及部分硬度压痕；
图5示出了T1和T2成分的预镀层钢板在热冲压后获得的热冲压成形构件的基体钢板的近界面区的脱碳程度(即HV 1/HV 2)与露点之间的关系；
图6示出了T1成分不同厚度的预镀层钢板经相同的脱碳处理和热冲压工艺后获得的热冲压成形构件的弯曲断裂应变及VDA弯曲角与厚度的关系；
图7示出了T1和T2成分1.4mm的预镀层钢板在热冲压后获得的热冲压成形构件的弯曲断裂应变与VDA弯曲角的关系；
图8示出了T1和T2成分1.4mm的预镀层钢板在热冲压后获得的热冲压成形构件的弯曲断裂应变随HV 1/HV 2比值的变化情况；
图9示出了T1和T2成分1.4mm的预镀层钢板在热冲压后获得的热冲压成形构件的抗拉强度随HV 1/HV 2比值的变化情况；及
图10示出了T1和T2成分1.4mm的预镀层钢板在热冲压后获得的热冲压成形构件的VDA峰值力随HV 1/HV 2比值的变化情况。
具体实施方式
要指出的是，初始低碳区对于韧性的改善是因为其消减或抵消了在热冲压成形过程中由于扩散导致的碳富集以及靠近基体钢板的软的铁素体相互扩散层的形成。因此，相比于现有的具有铝或铝合金镀层的预镀层钢板，在其他条件一致的情况下，在钢板表面保留一定厚度的初始低碳区将改善由镀层钢板制成的热冲压成形构件的韧性。但是应注意，为了避免提高韧性进行的脱碳导致VDA峰值力大幅下降，本发明人提出控制低碳区的厚度，即控制 脱碳程度，使其足以削弱或抵消碳富集。同时，本申请人也考虑了基体钢板的碳含量对后续热冲压加热过程中的碳富集存在的影响。热冲压加热过程中，当钢板具有相对较高的碳含量时，相应的碳富集也会严重，因此，期望提高预镀层钢板的基体钢板表面的脱碳程度，即，低碳区的碳含量不宜太高。相对地，当钢板具有较低的碳含量时，碳富集现象相对较弱，因此，脱碳程度可以适当降低，即表面低碳区的碳含量可以稍高一些。
鉴于上述，考虑到脱碳程度与露点的关系及基体钢板的碳含量，本发明提供了一种制造具有铝或铝合金预镀层的预镀层钢板的方法，其使得由预镀层钢板获得的热冲压成形构件具有优异的强度和韧性。根据本发明的涂镀方法包括：
a)脱碳处理：将基体钢板在H 2体积百分数为2～12％的H 2和N 2的环境气氛中加热至在740～880℃范围内，优选地在740～820℃范围内的第一温度并保温30～300s，其中，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内，对于0.10％≤C 0≤0.30％的情况，控制环境气氛的露点在-40～-15℃范围内，并且对于0.30％＜C 0≤0.50％的情况，控制环境气氛的露点在-36～-12℃范围内；
b)热浸镀：将加热后的基体钢板冷却到在610～680℃范围内的第二温度，之后浸入温度为610～680℃的镀液中进行热浸镀；
c)在基体钢板离开镀液后且在至少一个表面上的镀液凝固前，通过气刀吹扫来移除至少一个表面上多余的镀液以控制所述至少一个表面上的预镀层厚度w 1；及
d)将基体钢板冷却至室温以获得具有铝或铝合金预镀层的预镀层钢板，其中，预镀层厚度w 1为5～20μm，预镀层钢板的总厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm。
镀液组成包含以质量计：9～12％Si、4％以下的Fe、余量为Al以及不可避免的杂质。
此外，为进一步使得热冲压成形构件的韧性提升且保证高的VDA峰值力，优选地，对于0.10％≤C 0≤0.30％的情况，控制所述环境气氛的露点在-35～-17℃范围内，更优选地在-31～-19℃范围内。优选地，对于0.30％＜C 0≤0.50％的情况，控制所述环境气氛的露点在-30～-15℃范围内，更优选地在-27～-17℃范围内。
基于上述制造方法，本发明提供了一种涂覆有铝或铝合金预镀层的预镀层钢板，钢板总厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm，所述预镀层钢板包括基体钢板和在基体钢板的至少一个表面上的铝或铝合金的预镀层，
所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；
所述预镀层的预镀层厚度w 1为5～20μm，其中Al含量以质量计大于等于60％；
在所述基体钢板内邻接所述基体钢板与所述预镀层之间的界面存在初始低碳区，
(1)在0.10％≤C 0≤0.30％的情况下，
a)从所述界面至基体钢板内6μm处的碳含量C 1a满足53％C 0≤C 1a≤C 0；且
b)从所述界面至基体钢板内10μm处的碳含量C 1b满足75％C 0≤C 1b≤C 0且C 1b＞C 1a；
优选地，59％C 0≤C 1a≤90％C 0，同时，77.5％C 0≤C 1b≤95％C 0且C 1b＞C 1a；
更优选地，64％C 0≤C 1a≤82％C 0，同时，80％C 0＜C 1b≤91.5％C 0且C 1b＞C 1a；
(2)在0.30％＜C 0≤0.50％的情况下，
a)从所述界面至基体钢板内6μm处的碳含量C 2a满足42％C 0≤C 2a≤87％C 0；且
b)从所述界面至基体钢板内10μm处的碳含量C 2b满足65％C 0≤C 2b≤95％C 0且C 2b＞C 2a；
优选地，50％C 0≤C 2a≤75％C 0，同时，70％C 0≤C 2b≤86％C 0且C 2b＞C 2a；
更优选地，55％C 0≤C 2a≤70％C 0，同时，75％C 0＜C 2b≤85％C 0。
此外，本发明还提供了一种带铝或铝合金镀层的热冲压成形构件，热冲压成形构件的厚度为0.5～3.0mm，优选地0.7～2.3mm，更优选地0.8～2.0mm，由内至外所述热冲压成形构件包括：
基体钢板，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；及
铝或铝合金镀层，其厚度为10～26μm并且包括：邻接基体钢板的相互扩散层，该相互扩散层的厚度为6～14μm并且包括含Al的铁素体，其中Fe含量以质量计大于等于70％；及在相互扩散层外侧的Fe和Al的金属间化合物层；
对于0.10％≤C 0≤0.30％且抗拉强度在1300～1800MPa范围内的热冲压成形构件，从基体钢板与镀层之间的界面至基体钢板6μm内的近界面硬度HV 1为基体钢板的心部硬度HV 2的0.65～1.07倍且HV 2在400～550HV的范围内，对于0.30％＜C 0≤0.50％且抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.6～1.0倍且HV 2大于550HV。优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.283并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的98％；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.21并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的97％。
优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.70～1.0倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.65～0.90倍。 进一步优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.30并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.23并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的99％。
更优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.75～0.95倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.68～0.85倍。进一步优选地，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.31并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高至少2％；抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.24并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高。
HV 1和HV 2是使用5g的载荷力测得的维氏硬度值的10点平均值。
化学成分详细描述如下：
C：约0.10～0.50％
C是钢中最有效的固溶强化元素，为保证钢材的抗拉强度在1300MPa以上，需要碳含量大于等于约0.10％。但是，如果碳含量超过0.50％，其显微组织主要为脆性的高碳马氏体，其延性和韧性均较差，且抗氢脆性能显著下降。因此，本发明所用的高强度钢的C含量在约0.10～0.50％之间。
Mn：约0.50～10％
Mn是提高钢材淬透性以确保钢材强度的元素。当Mn含量低于约0.50％时，钢材的淬透性不足，难以获得高强度。但是，当Mn含量过高时，钢材容易出现中心带状偏析，对钢材的延性和韧性存在不利的影响，因此，本发明所用的高强度钢的Mn含量上限为约10％。
所获得的热冲压成形构件的韧性的改善主要来自于两个方面：
一、铝或铝合金镀层中的相互扩散层。相比于裸板，相互扩散层为软的高Al的铁素体，具有较好的韧性和塑性，该层的存在改善了热冲压成形构件的韧性。通过本发明的研究发现，相互扩散层厚度在6～14μm最佳。当扩散层厚度低于6μm时，基体钢板奥氏体化不充分，冷却后的组织均匀性较差，影响最终构件的综合性能。当扩散层厚度高于14μm时，说明热冲压加热时间过长或温度过高，将导致基体钢板奥氏体化过程中晶粒发生明显长大，恶化最终构件的韧性；
二、预镀层钢板的初始低碳区，其消减或抵消在热冲压过程中由于扩散的发生在最终镀层和基体钢板界面附近在基体钢板内发生堆积而形成的碳富集区，从而减少或避免在随后的冷却中碳富集区转变为脆性的高碳马氏体组织，改善了最终热冲压成形构件的韧性。
本发明采用检测热冲压成形构件的基体钢板的近界面硬度以及其与基体钢板的心部硬度之比的方式，来确定热冲压后构件的基体钢板的近界面的脱碳或碳富集程度。硬度比不宜过高，过高则说明基体钢板内仍存在较多的脆性高碳马氏体组织，即热冲压过程中的碳富集现象未能有效消减，对构件的韧性不利。但是也不宜过低，过低则说明脱碳严重，导致基体钢板表面硬度偏低且抗拉强度和VDA峰值力显著下降，即构件抵抗拉伸和弯曲变形失效的能力变差，影响最终构件的碰撞安全性和轻量化效果。故而，针对不同强度的热冲压成形构件，本申请要求在使用5g的载荷力进行维氏硬度测试时，从热冲压成形构件的基体钢板与镀层之间的界面至基体钢板内6μm内的近界面硬度HV 1与基体钢板中心区域的心部硬度HV 2具有以下关系：
对于0.10％≤C 0≤0.30％且抗拉强度在1300～1800MPa范围内的热冲压成形构件，从基体钢板与镀层之间的界面至基体钢板6μm内的近界面硬度HV 1为基体钢板的心部硬度HV 2的0.65～1.07倍且HV 2在400～550HV的范围内，对于0.30％＜C 0≤0.50％且抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.6～1.0倍且HV 2大于550HV。
优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.70～1.0倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.65～0.90倍。
更优选地，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.75～0.95倍；对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.68～0.85倍。
下面将参考示例性实施例来更详细地描述本发明。以下实施例或实验数据旨在示例性地说明本发明，本领域的技术人员应该清楚的是本发明不限于这些实施例或实验数据。
制备具有表1所示成分的基体钢板，相应的制造工艺如下：
a)炼钢：按照表1成分由真空感应炉、电炉或转炉冶炼，利用连铸技术生产铸坯，或直接采用薄板坯连铸连轧工艺；
b)热轧：将钢坯加热至1200℃保温2小时后在800～1200℃进行热轧，酸洗以清除热轧过程中产生的氧化皮，并在700℃以下进行卷曲，形成热轧钢卷；及
c)冷轧：将经过酸洗的热轧卷进行冷轧，冷轧压下量为30～70％，得到厚度为1.2mm、1.4mm和1.8mm的冷轧钢板。
表1基体钢板的化学成分(质量％，余量为Fe和其他不可避免杂质元素)
材料编号 C Mn Si Cr V B Al Ti Nb
T1 0.23 1.18 0.22 0.16 - 0.0025 0.034 0.04 -
T2 0.35 1.45 0.51 0.2 0.16 0.0019 0.56 - 0.04
其中，T1成分为22MnB5，采用该合金成分的热冲压成形构件的抗拉强度通常为1300～1800MPa。T2成分为35MnB5V，采用该合金成分的热冲压成形构件的抗拉强度可达1800MPa以上。文中关于化学元素含量的描述都是指质量百分数。
对上述冷轧钢板进行脱碳处理，例如能够在冷轧钢板的连续退火过程中进行，环境气氛为5％H 2+N 2(体积分数)。为得到不同的表面脱碳效果，对脱碳过程中的工艺参数，如加热温度、露点及保温时间，进行调整，参见表2。随后，将加热的钢板冷却至610～680℃范围内进行热镀，并且之后，利用气刀吹扫控制预镀层的厚度，从而得到带有不同厚度的铝或铝合金的预镀层的预镀层钢板。所述镀液的组成包含以质量计：9～12％Si；及余量为Al或Al合金以及不可避免的杂质。
表2表面脱碳处理工艺
Figure PCTCN2020124190-appb-000008
利用扫描电镜(SEM)测量预镀层钢板的预镀层厚度，并采用GDOES辉光放电发射光谱法测试从基体钢板与预镀层之间的界面至基体钢板内约18μm内的碳含量分布。
图1示出了示例A8的预镀层钢板的微观结构，其中，平均的预镀层厚度为约9.5μm。碳元素在该预镀层钢板的基体钢板中的分布结果在图2中示出。
图2示出了示例A8和B2预镀层钢板的基体钢板的碳元素分布情况。示例B2的曲线示出从所述界面至基体钢板内2μm处的碳含量已经达到基体钢板的碳含量C 0的约90％。相对地，示例A8的碳含量在2μm处仅为约60％C 0。在6μm处，示例B2的碳含量为约94.3％C 0，示例A8的碳含量为约69.1％C 0。在10μm处，示例B2的碳含量为约 96.8％C 0，示例A8的碳含量为约83.8％C 0。示例B2的碳含量比示例A8的高，这是因为示例A8的露点比示例B2的高，所以脱碳程度增大，使得示例A8的初始低碳区厚度比示例B2厚。
各示例的预镀层厚度及初始低碳区的碳含量测量结果见表3。
表3预镀层厚度及初始低碳区的碳含量(C 0为基体钢板的碳含量)
Figure PCTCN2020124190-appb-000009
由表2和表3可知，对于T1成分的示例A0～A6，露点由-38℃升至-22℃，从所述界面至基体钢板内6μm处的碳含量在70.9～95.5％C 0范围内，10μm处的碳含量比在6μm处高且在84.5～98.7％C 0范围内。相对地，示例B0的露点最低，达到-41℃，因此，脱碳程度最轻，从所述界面至基体钢板内6μm和10μm处的碳含量分别为基体钢板的约97％C 0和约99％C 0。而示例B1的露点最高，达到-5℃，因此，脱碳程度最高，从所述界面至基体钢板内6μm和10μm处的碳含量分别不足基体钢板的约27％C 0和约59％C 0。
对于T2成分的示例A7～A11，露点由-31℃升至-17℃，从所述界面至基体钢板内6μm处的碳含量在55.3～76.4％C 0范围内，10μm处的碳含量比在6μm处高且在72.6～87.2％C 0范围内。相对地，示例B2的露点为-39℃，基体钢板表面轻微脱碳，从所述界面至基体钢板内6μm和10μm处的碳含量分别为约94％C 0和约97％C 0。同样，示例B3的露点为-5℃，基体钢板表面脱碳严重，从所述界面至基体钢板内6μm和10μm处的碳含量 分别不足基体钢板的约25％C 0和约55％C 0。
图3示出了T1和T2成分的预镀层钢板的初始低碳区的碳含量与露点之间的关系。由图可知，随着露点的升高，初始低碳区的碳含量降低，相应地，低碳区厚度增加，即，脱碳程度增加。同时，与T1成分相比，由于T2成分的碳含量更高，所以在相同的露点条件下，T2成分示例的脱碳效果更加明显，故而T2成分示例的初始低碳区的碳含量比T1成分低。
对上述各示例预镀层钢板进行平板热冲压模拟，热冲压工艺为：
1)将预镀层钢板加热至930℃保温320s，随后冷却在700℃以上进行热冲压，保压时间6～10s，并在模具内冷却至100℃以下取出；
2)为模拟汽车零件实际的使用情况，待钢板冷至室温，放入170℃的回火炉中保温20min后以模拟零件涂装烘烤过程，随后，取出空冷至室温。
由预镀层钢板获得的热冲压成形构件包括基体钢板和其外侧的镀层。对热冲压态的构件测试基体钢板的近界面区和中心区的维氏硬度，并对回火态的构件进行室温拉伸性能和VDA弯曲性能的检测。
维氏硬度测试方法如下：选择5g力的压头载荷F HV，分别在从基体钢板与镀层之间的界面至基体钢板内6μm内和基体钢板的心部压入试样表面，随后卸除试验力并在SEM下测量压痕的对角线长度d 1和d 2。利用下述等式(8)计算相应的从基体钢板与镀层之间的界面至基体钢板内6μm内的近界面硬度和基体钢板的心部硬度：
Figure PCTCN2020124190-appb-000010
为减小硬度测量误差，所有硬度最终结果均采用测试10点取平均值的方式。图4示出了T2成分示例A8热冲压后的局部微观结构以及部分硬度压痕，其中，镀层厚度约为14.7～16.9μm，相互扩散层厚度约为6.7～7.5μm。
室温拉伸性能和VDA弯曲性能的测试方法分别参考GB/T 228.1和VDA 238-100标准，试样的取样方向为轧制方向。同样，为减小测量误差，选取三组试样进行拉伸性能和VDA弯曲性能测试，最终结果均为三组测试结果的平均值。
弯曲断裂应变测试方法如下：(1)利用静态三点弯曲实验确定试样的VDA弯曲角，α peak；(2)基于实验结果，选取至少三组中断弯曲角α L(即试样在承载状态下的弯曲角)进行中断弯曲实验，保证α L≥50％α peak；(3)当试样弯曲至α L时停止加载，测量试样在卸载状态下的弯曲角α UL；(4)将卸载试样放置于光学显微镜下，测量最严重变形区的内、外 表面半径R i和R o；(5)根据等式(9)计算出不同α UL的卸载状态下的试样最严重变形区外表面等效(塑性)应变ε，即，试样弯曲至α L时试样最严重变形区的外表面的等效应变，从而建立ε-α L关系；(6)根据拟合结果，利用外推方法得出试样的弯曲断裂应变(即α L等于α peak时的ε)，
Figure PCTCN2020124190-appb-000011
维氏硬度、拉伸性能和VDA弯曲性能的最终结果见表4。
表4热冲压成型构件的相关测量参数
Figure PCTCN2020124190-appb-000012
由于基体钢板和预镀层之间的扩散，所以在热冲压成形过程之后，初始低碳区缩窄，甚至不复存在。这使得难以对低碳区的厚度进行测量，故而本申请利用热冲压成形构件的基体钢板的近界面区的硬度与心部硬度比值(即HV 1/HV 2)来反映脱碳程度。
图5示出了T1和T2成分的预镀层钢板在热冲压后获得的热冲压成形构件的基体钢板的近界面区的脱碳程度(即HV 1/HV 2)与露点之间的关系。由图可知，随着露点的升高，HV 1/HV 2的比值呈下降趋势，这与图3所示的碳含量随露点升高而下降一致。HV 1/HV 2的比值与预镀层钢板的脱碳程度是相对应的。
注意到，对于T1成分，虽然示例A0的近界面硬度与心部硬度比为1.05，显示出仍存在碳富集，但是VDA弯曲角和弯曲断裂应变已经分别提高至62.8°和0.294。这是因为 虽然示例A0的露点较低，使得其基体钢板表面仅轻微脱碳，但是该轻微脱碳已能够消减热冲压过程中由于相互扩散导致的碳富集现象，使得VDA弯曲性能相对于现有技术进一步提高。
厚度同样为1.4mm的示例A1～A4的近界面硬度与心部硬度之比在0.81～0.95范围内，这意味着经热冲压后，在基体钢板与镀层之间的界面附近无碳富集现象。这是因为与示例A0相比，示例A1～A4的初始脱碳区碳含量降低，即脱碳程度相对增加，进一步抵消了A0中存在的轻微碳富集现象。由于无碳富集现象，故而在热冲压成形之后的冷却过程中，近界面处基本没有生成脆性的高碳马氏体。因此，相较于示例A0，示例A1～A4热冲压后的韧性得到进一步地改善，VDA弯曲角可达64.2～66.4°，弯曲断裂应变在0.309～0.315。由此可见，使预镀层钢板的基体表面存在一定程度的脱碳，能够有效消减热冲压过程中由于扩散导致的碳富集，从而减少脆性马氏体的生成，改善韧性，表现为VDA弯曲角和弯曲断裂应变的提高。相对的，示例B0的近界面硬度与心部硬度比为1.12，碳富集现象明显，这是因为偏低的露点使得基体钢板表面几乎没有脱碳，从而无法有效消减碳富集现象，因此，VDA弯曲性能无法提升，VDA弯曲角仅为56.7°，弯曲断裂应变为0.262。
对于T2成分，由于其基体钢板的碳含量较高，所以其对热冲压过程中碳富集现象更加敏感。示例B2的近界面硬度与心部硬度比为1.03，显示出轻微的碳富集。但是由于T2成分的碳含量较高，因此，对于示例B2，即使轻微的碳富集也不利于热冲压成形构件的韧性的改善。结果也是如此，示例B2的VDA弯曲角仅为45.4°，弯曲断裂应变约0.201。为了进一步抵消碳富集，对于T2成分，需适当提高预涂镀过程中基体钢板表面的脱碳程度。示例A7～A11的近界面硬度与心部硬度之比在0.70～0.93范围内，相应地VDA弯曲角提高至50.3～57.7°，弯曲断裂应变也提高至0.224～0.272，这说明碳富集现象对热冲压成形构件的韧性的影响已经被消除。
图6示出了由不同厚度的T1成分的预镀层钢板经相同的脱碳处理和热冲压工艺后获得的热冲压成形构件的弯曲断裂应变及VDA弯曲角与构件的厚度的关系。由图可知，随着热冲压成形构件的厚度增大，VDA弯曲角大致线性降低，但弯曲断裂应变保持大致恒定。这是因为，对于确定的材料成分，理论上相同的处理工艺将导致所获的热冲压成形构件的近表层状态相同，而VDA弯曲实验主要导致的就是构件最外表面(即近表层)发生断裂，因此，弯曲断裂应变不随构件厚度变化而变化。故而，用弯曲断裂应变来评价不同厚度的材料的碰撞安全性是可靠的。
图7示出了由1.4mm的T1和T2成分的预镀层钢板在热冲压后获得的热冲压成形构 件的弯曲断裂应变与VDA弯曲角的关系。在同一厚度下，两种成分的热冲压成形构件的弯曲断裂应变ε与VDA弯曲角α peak均呈线性关系，满足下述等式(10)：
ε＝0.00554α peak-0.05    (10)
根据该等式，对于T1成分，当α peak为60°时，ε为0.2824，优选地，当α peak为63.2°时，ε约为0.30，更优选地，当α peak为65°时，ε约为0.31时。对于T2成分，当α peak为47°时，ε约为0.21，优选地，当α peak为50.5°时，ε约为0.23时，更优选地，当α peak为52.3°时，ε约为0.24。相比于T1成分的热冲压成形构件，强度高达1900MPa的T2成分的热冲压成形构件的VDA弯曲角较低，弯曲断裂应变也较低，即强度的提升伴随着构件韧性的下降。
图8示出了由1.4mm的T1和T2成分的预镀层钢板在热冲压后获得的热冲压成形构件的弯曲断裂应变随HV 1/HV 2比值的变化情况。如图所示，对于两种成分，随着HV 1/HV 2的比值降低，即近界面区的脱碳程度增加，热冲压成形构件的弯曲断裂应变均呈上升的趋势，但趋势逐渐减缓。这是由于弯曲断裂应变仅与基体钢板的近表层状态有关，对其进行脱碳处理可显著改善近表层的韧性和塑性。对于带铝或铝合金预镀层的预镀层钢板，在热冲压后，由于基体钢板表面形成了一定厚度的高Al的铁素体的相互扩散层，因此，在基体钢板的近界面区不存在严重的碳富集时，钢板将具有良好的韧性，即具有高的弯曲断裂应变。而进一步提高近界面区的脱碳程度对近表层的韧性改善的贡献较小，使得弯曲断裂应变的增加有限。
图9和图10示出了由1.4mm的T1和T2成分的预镀层钢板在热冲压后获得的热冲压成形构件的抗拉强度和VDA峰值力随HV 1/HV 2比值的变化情况。注意到，随着HV 1/HV 2的比值降低，即近界面区的脱碳程度增加，热冲压成形构件的抗拉强度出现小幅下降。例如，对于T1成分，以示例B0(露点为-41℃，由前述数据可知，接近无脱碳)为基准，脱碳程度最高的示例B1(露点为-5℃)的HV 1/HV 2的比值不足0.5，其抗拉强度相比于示例B0降低了约61MPa，降幅约为3.94％。相对的，脱碳程度比示例B1低的示例A0-A4的抗拉强度相对于示例B0的降幅在1％以内，如此小的降幅可忽略不计。类似地，对于T2成分，以示例B2(露点为-39℃)为基准，脱碳程度最高的示例B3(露点为-5℃)的HV 1/HV 2的值不足0.5，其抗拉强度相比于示例B2降低了约67MPa，降幅约3.48％。相对的，脱碳程度比示例B3低的示例A7-A11的抗拉强度相对于示例B0的降幅在2％以内，如此小的降幅可忽略不计。
相对的，随着HV 1/HV 2的比值降低(即，脱碳区厚度增加)，热冲压成形构件的 VDA峰值力虽然也出现下降趋势，但是在此之前经历了升高阶段。详细解释如下。对于T1成分，以示例B0(露点为-41℃，由前述数据可知，接近无脱碳)为基准，脱碳程度较低的示例A0-A4(HV 1/HV 2比值在0.81～1.05范围内)的VDA峰值力相对于示例B0没有下降，反而升高了。这是因为示例B0(HV 1/HV 2比值为1.12)存在明显碳富集现象，由于碳富集形成的脆性高碳马氏体组织在较低载荷下就能够引起开裂，所以试样早期就发生失效。实际上，示例B0相对于完全无脱碳已经消除了些许碳富集，也就是说，在完全无脱碳的情况下，试样的VDA峰值力应该会更低。相对的，示例A0的碳富集被部分抵消，使得碳富集现象轻微，其VDA峰值力相对于示例B0的峰值力升高。示例A1中的碳富集被完全抵消，故而相对于示例A0的VDA峰值力进一步升高。其他示例A2-A4相对于示例A1进一步脱碳，然而VDA峰值力相对于示例A1逐渐下降，这是因为在碳富集已经完全被抵消的情况下，进一步增加脱碳将使峰值力降低。故而，当脱碳程度大到使脱碳导致的VDA峰值力降低大于消除碳富集带来的VDA峰值力提高时，将表现为VDA峰值力比无脱碳情况下低。例如，示例B1(露点为-5℃)的VDA峰值力相比于示例B0明显下降，降低了约500N，降幅约5.52％。这是不期望的。
类似地，对于T2成分，以示例B2(露点为-39℃)为基准，脱碳程度较低的示例A7-A11(HV 1/HV 2比值在0.7～0.93范围内)的VDA峰值力相对于示例B0同样升高了。然而，当脱碳程度大到使脱碳导致的VDA峰值力降低大于消除碳富集带来的VDA峰值力提高时，VDA峰值力将比无脱碳低。例如，示例B3(露点为-5℃)的VDA峰值力相比于示例B2明显下降，降低了约936N，降幅约7.74％。
上述实验数据表明，在碳富集未完全消除的情况下，增加脱碳程度将使VDA弯曲力提高。此外与前述理论上一致，在碳富集消除的情况下，VDA峰值力将随着低碳区厚度的增加下降。但是在脱碳带来VDA峰值力下降不大于抵消碳富集带来的VDA峰值力提高的情况下，通过脱碳处理仍能够提高VDA峰值力，现有技术完全未意识到这一点。
考虑到前述，由于过度增加预镀层钢板的基体钢板的表面脱碳程度将会降低热冲压成形构件的抗拉强度和VDA峰值力，特别是VDA峰值力，其降幅更加显著，这将影响热冲压成形构件的碰撞安全性和轻量化效果，所以，在评价热冲压成形构件的碰撞安全性的情况下，必须要将VDA峰值力作为重要参数考虑进去。为此，本申请提出控制初始脱碳程度，使得其既能够消减或消除脆性高碳马氏体组织，又能保证脱碳带来VDA峰值力下降不过度大于消除碳富集带来的VDA峰值力提高，从而改善热冲压成形构件的碰撞安全性和轻量化效果。
为了保证碰撞安全性，构件在具有高韧性的同时仍需要具有高的VDA峰值力。由图10可知，对于拉伸强度为约1500MPa的1.4mm的T1成分的热冲压成形构件，VDA峰值力不小于9000N(不小于接近无脱碳的示例B0的VDA峰值力的98％)，则HV 1/HV 2比值应不小于0.65。此外，结合图7-9及表4的数据可知，对于拉伸强度为约1500MPa的1.4mm的T1成分的热冲压成形构件，为使其VDA弯曲角大于60°，弯曲断裂应变应不小于0.283，对应地，HV 1/HV 2比值不大于1.07。故而，为了使热冲压成形构件的VDA峰值力不低于9000N且VDA弯曲角大于60°，本申请要求HV 1/HV 2在约0.65～1.07的范围内。在这种情况下，结合图5，对应地本申请要求露点在约-40～-15℃的范围内以满足热冲压成形构件的期望性能。相应地，结合图3，预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 1a满足53％C 0≤C 1a≤C 0，且10μm处的碳含量C 1b满足75％C 0≤C 1b≤C 0且C 1b＞C 1a。
优选地，本申请规定对于T1成分，VDA峰值力不小于9150N(比接近无脱碳的示例B0的VDA峰值力高约0.84％)，故HV 1/HV 2比值应不小于0.70。同时，期望VDA弯曲角不小于63.2°，则弯曲断裂应变应不小于0.30，对应地，HV 1/HV 2比值不大于1.00。故而，为了使热冲压成形构件的VDA峰值力不小于9150N且VDA弯曲角不小于63.2°，HV 1/HV 2在约0.7～1.0的范围内，对应地，露点在约-35～-17℃的范围，并且预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 1a满足59％C 0≤C 1a≤90％C 0，且10μm处的碳含量C 1b满足77.5％C 0≤C 1b≤95％C 0且C 1b＞C 1a。
更优选地，本申请规定对于T1成分，VDA峰值力不小于9300N(比接近无脱碳的示例B0的VDA峰值力高约2.5％)，故而HV 1/HV 2比值应不小于0.75。同时，期望VDA弯曲角不小于65°时，则弯曲断裂应变应不小于0.31，对应地，HV 1/HV 2比值不大于0.95。故而，为了使热冲压成形构件的VDA峰值力不小于9300N且VDA弯曲角不小于65°，HV 1/HV 2在约0.75～0.95的范围内，对应地，露点在约-31～-19℃的范围，并且预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 1a满足64％C 0≤C 1a≤82％C 0，且10μm处的碳含量C 1b满足80％C 0＜C 1b≤91.5％C 0且C 1b＞C 1a。
类似地，对于拉伸强度为约1900MPa的1.4mm的T2成分的热冲压成形构件，VDA峰值力不低于11800N(不小于接近无脱碳的示例B2的VDA峰值力的97％)，故而HV 1/HV 2比值应不小于0.6。同时期望VDA弯曲角不小于47°，此时，弯曲断裂应变应至少不小于0.21，对应地，HV 1/HV 2比值不大于1。故而，HV 1/HV 2在约0.6～1.0的范围内，结合图5，本申请要求露点在约-36～-12℃的范围内以满足热冲压成形构件的期望性能。相应 地，结合图3，预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 2a满足42％C 0≤C 2a＜87％C 0，且10μm处的碳含量C 2b满足65％C 0≤C 2b≤95％C 0且C 2b＞C 2a。
优选地，对于T2成分，VDA峰值力不低于12000N(不小于接近无脱碳的示例B2的VDA峰值力的99％)且VDA弯曲角不小于50.5°，对应地，弯曲断裂应变应不小于0.23，HV 1/HV 2比值在约0.65～0.90的范围内。在这种情况下，本申请要求露点在约-30～-15℃的范围内以满足热冲压成形构件的期望性能。相应地，预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 2a满足50％C 0≤C 2a≤75％C 0，且10μm处的碳含量C 2b满足70％C 0≤C 2b≤86％C 0且C 2b＞C 2a。
更优选地，对于T2成分，VDA峰值力不低于12200N(比接近无脱碳的示例B2的VDA峰值力高约0.93％)且VDA弯曲角不小于52.3°，对应地，弯曲断裂应变应不小于0.24，HV 1/HV 2比值在约0.68～0.85的范围内。在这种情况下，本申请要求露点在约-27～-17℃的范围内以满足热冲压成形构件的期望性能。相应地，预镀层钢板的碳含量应满足：从所述界面至基体钢板内6μm处的碳含量C 2a满足55％C 0≤C 2a≤70％C 0，且10μm处的碳含量C 2b满足75％C 0＜C 2b≤85％C 0。
最后，为保证预镀层钢板基体表面具有高的脱碳程度，示例B1和B3的露点已经提高至-5℃，实际生产过程中，炉内氧化气氛严重，基体钢板表面已出现显著的氧化色，在后续涂镀时出现了漏镀、镀层结合力差等问题，生产管控困难，生产成本也随之提高。
考虑到上述，本发明通过设置适当的脱碳工艺控制预镀层钢板的基体钢板的初始低碳区的碳含量(即脱碳程度)，从而使得根据本发明获得的热冲压成形构件不仅具有提高的韧性，而且具有高的抗拉强度和VDA峰值力，改善了热冲压成形构件的碰撞安全性。
以上实施例和实验数据旨在示例性地说明本发明，本领域的技术人员应该清楚的是本发明不仅限于这些实施例，在不脱离本发明保护范围的情况下，可以进行各种变更。

Claims (16)

  1. 一种涂覆有铝或铝合金预镀层的预镀层钢板，所述预镀层钢板的总厚度为0.5～3.0mm，所述预镀层钢板包括基体钢板和在基体钢板的表面上的铝或铝合金的预镀层，
    所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；
    所述预镀层的预镀层厚度w 1为5～20μm，其中Al含量以质量计大于等于60％；
    在所述基体钢板内邻接所述基体钢板与所述预镀层之间的界面存在初始低碳区，
    (1)在0.10％≤C 0≤0.30％的情况下，
    a)从所述界面至基体钢板内6μm处的碳含量C 1a满足53％C 0≤C 1a≤C 0；且
    b)从所述界面至基体钢板内10μm处的碳含量C 1b满足75％C 0≤C 1b≤C 0且C 1b＞C 1a；
    (2)在0.30％＜C 0≤0.50％的情况下，
    a)从所述界面至基体钢板内6μm处的碳含量C 2a满足42％C 0≤C 2a≤87％C 0；且
    b)从所述界面至基体钢板内10μm处的碳含量C 2b满足65％C 0≤C 2b≤95％C 0且C 2b＞C 2a。
  2. 根据权利要求1所述的预镀层钢板，其中，所述总厚度为0.7～2.3mm，其中，
    59％C 0≤C 1a≤90％C 0，同时77.5％C 0≤C 1b≤95％C 0且C 1b＞C 1a；
    50％C 0≤C 2a≤75％C 0，同时70％C 0≤C 2b≤86％C 0且C 2b＞C 2a。
  3. 根据权利要求1所述的预镀层钢板，其中，所述总厚度为0.8～2.0mm，其中，
    64％C 0≤C 1a≤82％C 0，同时80％C 0＜C 1b≤91.5％C 0且C 1b＞C 1a；
    55％C 0≤C 2a≤70％C 0，同时75％C 0＜C 2b≤85％C 0。
  4. 根据权利要求1-3中任一项所述的预镀层钢板，其中，所述基体钢板以质量百分比计包含以下成分：0.10～0.50％的C，0.50～10％的Mn，0～0.01％的B，0～0.4％的Nb+Ti+V，0.01～2％的Si，0.01～2％的Al，0.01～5％的Cr+Ni+Mo+Cu且0～2％的Cr、0～2％的Ni、0～2％的Mo及0～2％的Cu，以及余量为Fe及不可避免的杂质元素。
  5. 一种制造具有铝或铝合金预镀层的预镀层钢板的方法，包括：
    a)脱碳处理：将基体钢板在H 2体积百分数为2～12％的H 2和N 2的环境气氛中加热至在740～880℃范围内的第一温度并保温30～300s，其中，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内，对于0.10％≤C 0≤0.30％的情况，控制所述环境气氛的露点在-40～-15℃范围内；对于0.30％＜C 0≤0.50％的情况，控制所述环境气氛的露点在-36～-12℃范围内；
    b)热浸镀：将加热后的基体钢板冷却到在610～680℃范围内的第二温度，之后浸入温度为610～680℃的镀液中进行热浸镀；
    c)在基体钢板离开镀液后且在表面上的镀液凝固前，通过气刀吹扫来移除表面上多余的镀 液以控制表面上的预镀层厚度w 1；及
    d)将基体钢板冷却至室温以获得具有铝或铝合金预镀层的预镀层钢板，
    所述预镀层厚度w 1为5～20μm，所述预镀层钢板的总厚度为0.5～3.0mm。
  6. 根据权利要求5所述的方法，其中，所述第一温度在740～820℃范围内，所述总厚度为0.7～2.3mm；对于0.10％≤C 0≤0.30％的情况，控制所述环境气氛的露点在-35～-17℃范围内；对于0.30％＜C 0≤0.50％的情况，控制所述环境气氛的露点在-30～-15℃范围内。
  7. 根据权利要求5所述的方法，其中，所述总厚度为0.8～2.0mm；对于0.10％≤C 0≤0.30％的情况，控制所述环境气氛的露点在-31～-19℃范围内；对于0.30％＜C 0≤0.50％的情况，控制所述环境气氛的露点在-27～-17℃范围内。
  8. 一种带铝或铝合金镀层的热冲压成形构件，所述热冲压成形构件的厚度为0.5～3.0mm，由内至外所述热冲压成形构件包括：
    基体钢板，所述基体钢板的碳含量C 0在0.10～0.50％的范围内，锰含量在0.50～10％的范围内；及
    铝或铝合金镀层，其厚度为10～26μm并且包括：邻接基体钢板的相互扩散层，所述相互扩散层的厚度为6～14μm并且包括含Al的铁素体，其中Fe含量以质量计大于等于70％；及在所述相互扩散层外侧的Fe和Al的金属间化合物层；
    对于0.10％≤C 0≤0.30％且抗拉强度在1300～1800MPa范围内的热冲压成形构件，从基体钢板与镀层之间的界面至基体钢板6μm内的近界面硬度HV 1为基体钢板的心部硬度HV 2的0.65～1.07倍且HV 2在400～550HV的范围内；
    对于0.30％＜C 0≤0.50％且抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.6～1.0倍且HV 2大于550HV。
  9. 根据权利要求8所述的热冲压成形构件，其中，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.283并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的98％；
    抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.21并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的97％。
  10. 根据权利要求8所述的热冲压成形构件，其中，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.70～1.0倍；
    对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.65～0.90倍。
  11. 根据权利要求10所述的热冲压成形构件，其中，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.30并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高；
    抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.23并且VDA峰值力不小于由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力的99％。
  12. 根据权利要求8所述的热冲压成形构件，其中，对于抗拉强度在1300～1800MPa范围内的热冲压成形构件，HV 1为HV 2的0.75～0.95倍；
    对于抗拉强度高于1800MPa的热冲压成形构件，HV 1为HV 2的0.68～0.85倍。
  13. 根据权利要求12所述的热冲压成形构件，其中，抗拉强度在1300～1800MPa范围内的热冲压成形构件的弯曲断裂应变不小于0.31并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高至少2％；
    抗拉强度高于1800MPa的热冲压成形构件的弯曲断裂应变不小于0.24并且VDA峰值力比由具有相同成分且经历相同热冲压过程的无脱碳预镀层钢板获得的热冲压成形构件的峰值力高。
  14. 根据权利要求8或9所述的热冲压成形构件，其中，对于1.4mm的抗拉强度为1500MPa的热冲压成形构件，VDA弯曲角不小于60°且VDA峰值力不小于9.0kN；对于1.4mm的抗拉强度为1900MPa的热冲压成形构件，VDA弯曲角不小于47°且VDA峰值力不小于11.8kN。
  15. 根据权利要求10或11所述的热冲压成形构件，其中，对于1.4mm的抗拉强度为1500MPa的热冲压成形构件，VDA弯曲角不小于63.2°且VDA峰值力不小于9.15kN；对于1.4mm的抗拉强度为1900MPa的热冲压成形构件，VDA弯曲角不小于50.5°且VDA峰值力不小于12.0kN。
  16. 根据权利要求12或13所述的热冲压成形构件，其中，对于1.4mm的抗拉强度为1500MPa的热冲压成形构件，VDA弯曲角不小于65°且VDA峰值力不小于9.3kN；对于1.4mm的抗拉强度为1900MPa的热冲压成形构件，VDA弯曲角不小于52.3°且VDA峰值力不小于12.2kN。
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