WO2024104280A1 - 一种低碳的高韧性热冲压成形构件及钢板 - Google Patents

一种低碳的高韧性热冲压成形构件及钢板 Download PDF

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WO2024104280A1
WO2024104280A1 PCT/CN2023/131148 CN2023131148W WO2024104280A1 WO 2024104280 A1 WO2024104280 A1 WO 2024104280A1 CN 2023131148 W CN2023131148 W CN 2023131148W WO 2024104280 A1 WO2024104280 A1 WO 2024104280A1
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hot stamping
steel sheet
less
component
steel
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PCT/CN2023/131148
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English (en)
French (fr)
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易红亮
陈仪霖
冯婷婷
周澍
杨达朋
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育材堂（苏州）材料科技有限公司
东北大学
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Publication of WO2024104280A1 publication Critical patent/WO2024104280A1/zh

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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
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • 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
    • 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
    • B21D37/00—Tools as parts of machines covered by this subclass
    • B21D37/16—Heating or cooling
    • 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
    • C21D7/00—Modifying the physical properties of iron or steel by deformation
    • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
    • 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
    • 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
    • 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
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • 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
    • C22C38/06—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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • 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
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • 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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese

Definitions

  • the invention relates to a low-carbon high-toughness hot stamping formed component and a steel plate.
  • the room temperature tensile test (GB/T 228.1 standard) is a common method for measuring the strength of materials.
  • the yield strength and tensile strength obtained reflect the ability of the material to resist tensile deformation failure.
  • the more commonly used method for testing the toughness of materials is the static three-point bending test (ie VDA bending test, VDA 238-100 standard). The test can effectively measure the fracture strain of the material, thereby reflecting the ability of the material to resist bending deformation failure.
  • parts in the aforementioned specific areas are generally made of hot stamping steel 22MnB5. Although this steel plate has a tensile strength of more than 1400MPa, its toughness is low, and it is difficult to meet the increasingly stringent collision regulations and deformation energy absorption requirements, and there are safety hazards.
  • CN107810281A provides a press-hardened steel and a press-hardened component made of such steel.
  • the press-hardened steel component is mainly used for structural elements in motor vehicles for anti-intrusion or energy absorption functions. Satisfactory tensile strength, bendability and weld toughness are obtained by making the carbon content in the steel meet 0.062% ⁇ C ⁇ 0.095%, and the content ratio of martensite and bainite for improving ductility in the microstructure of the hot stamping component is ensured by making the content of hardenability elements C, Mn, Si, and Cr in the steel meet 1.5% ⁇ (C+Mn+Si+Cr) ⁇ 2.7%.
  • the patent teaches to use more bainite to improve the toughness and bendability of the component.
  • the patent technology specifically requires that the content of Cr is not higher than 0.1%, because the addition of Cr is not conducive to the formation of bainite for improving ductility.
  • the content of Cr is not higher than 0.1%, because the addition of Cr is not conducive to the formation of bainite for improving ductility.
  • CN104769138A found that forming a decarburized zone with a depth p50% of 6 to 30 ⁇ m on the surface of the base steel plate before coating helps the final hot-formed part to obtain high bendability, wherein the depth p50% is The carbon content is equal to the depth at which the carbon content of the base steel sheet is 50%.
  • the VDA bending angle of 1.8mm thick galvanized 22MnB5 hot-formed steel sheet can reach more than 55°.
  • the decarburized area of the above-mentioned specific thickness depends largely on the annealing conditions before coating, which increases the difficulty of production.
  • the patent ignores the adverse effects of the decarburized layer on the strength, surface hardness, peak load during bending, and subsequent coating quality of the material.
  • CN1717499A provides a high-strength steel plate for cold forming and a manufacturing method.
  • the composition of the steel plate is controlled to be C: 0.05-0.15%, Si: 0.3-2.0%, Mn: 1.0-2.8%, Al: 0.005-0.5%, N: 0.006% or less, and the rest is Fe and unavoidable impurities, and Mn/C ⁇ 12 and Si/C ⁇ 4.
  • the ratio between C, Mn and Si is designed to delay the phase transformation of pearlite and cementite during the alloying treatment after coating.
  • the microstructure of the steel plate of this patent is mainly ferrite and contains 3-20% martensite and residual austenite.
  • the present invention is made in view of the above problems existing in the prior art.
  • the steel sheet matrix of the hot stamping steel sheet of the present invention comprises, by mass percentage: 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, Fe ⁇ 95% and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy Mn+0.26Si+1.3Cr ⁇ 2.20%.
  • the steel plate matrix further comprises, in mass percentage, at least one of: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01%, and 0.0001% ⁇ REM ⁇ 0.01%.
  • the steel plate matrix of the steel plate for hot stamping forming of the present invention comprises, by mass percentage, the following: 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, and the balance is Fe and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy Mn+0.26Si+1.3Cr ⁇ 2.20%.
  • the steel plate matrix comprises, by mass percentage, 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, and at least one of the following: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.0 1% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01% and 0.0001% ⁇ REM ⁇ 0.01%, and 0.0001% ⁇ W+Mo+Ni+Cu+Co+Sn+Sb+Ca+
  • the invention obtains improved hardenability by rationally matching elements such as Mn, Cr, Si, etc. under the premise of low C content.
  • Ms 520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu.
  • 350°C ⁇ Ms ⁇ 410°C is satisfied, and more preferably, 360°C ⁇ Ms ⁇ 400°C.
  • the calculation formula of Ms is obtained from "Steel Heat Treatment Handbook" by A.V.Sverdlin and A.R.Ness (G.E.Totten and M.A.H.Howes, Ed., Marcel Dekker Inc., New York, 1997).
  • the calculation formula of Ac3 is from Zeng Qiying's "Quantitative Relationship of the Effect of Alloy Elements in Steel on Phase Transformation Points Ac1 and Ac3 " (Physical and Chemical Testing - Physical Section, 1985, 18(5):47-49).
  • the hot stamping steel sheet may have a metal coating, and preferably, the metal coating may be an aluminum-based alloy coating or a zinc-based alloy coating.
  • the steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage: 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb + V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.1 00%, 0.0001% ⁇ S ⁇ 0.100%, Fe ⁇ 95% and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy Mn+0.26Si+1.3Cr ⁇ 2.20%; wherein, in terms of area percentage, the microstructure of the steel plate matrix comprises: less than 5% of bainite, less than 3% of austenite, less than 3% of ferrite and less than 0.2% of
  • the steel plate matrix of the hot stamping formed component further contains, in mass percentage, at least one of the following: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01%, and 0.0001% ⁇ REM ⁇ 0.01%.
  • the steel plate matrix of the hot stamping formed component of the present invention contains, in mass percentage, the following: 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, and the balance is Fe and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy Mn+0.26Si+1.3Cr ⁇ 2.20%.
  • the steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage: 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, and at least one of the following: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30% , 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01% and 0.0001% ⁇ REM ⁇ 0.01%, of which 0.0001% ⁇ W+Mo+Ni+Cu
  • the hot stamping formed components of the present invention are optimized by reasonably matching various alloy elements such as Mn, Cr, Si, etc., so that high strength can be achieved while ensuring high toughness, meeting the requirements of collision energy absorption and lightweight in specific areas of the car body.
  • the sum of bainite, ferrite and retained austenite does not exceed 3%.
  • the component may have a metal coating.
  • the metal coating may be an alloy coating based on aluminum or an alloy coating based on zinc.
  • the component has a yield strength YS of 880-1100 MPa and a tensile strength TS of not less than 1000 MPa, an elongation of not less than 5%, and a fracture strain ⁇ of not less than 0.6.
  • the component has a yield strength of 930-1050 MPa and a tensile strength of not less than 1060 MPa, an elongation of not less than 7%, and a fracture strain of not less than 0.7.
  • the bending energy absorption W of the static three-point bending of the component and the plate thickness t satisfy: W/t 2 ⁇ 2.88 ⁇ 10 4 N/mm.
  • Another object of the present invention is a hot stamping method for producing the above-mentioned hot stamping formed component, which comprises:
  • Hot stamping process After completing step B, the heated blank or preform is transferred to a press for hot deformation and cooled to below 300°C in the die to obtain a hot stamping formed component, wherein the average cooling rate of the hot stamping formed component from 700°C to the Ms temperature in the area where no hot deformation occurs is not less than 40°C/s, and the hot deformation occurs in the area at a higher average cooling rate.
  • the Ms temperature of the hot stamping formed component satisfies: 350°C ⁇ Ms ⁇ 410°C.
  • step B is completed in a heating furnace, the maximum temperature in the furnace is in the range of 900-980°C, and the time (minutes, including heating and insulation time) of the blank or preform in the furnace satisfies: the thickness of the blank or preform is t (mm) ⁇ (t ⁇ 3+10).
  • step C the thermal deformation of at least one region of the blank or preform does not exceed 10%, and the average cooling rate of the at least one region from 700° C. to the Ms temperature is not less than 50° C./s.
  • step C the thermal deformation of at least one region of the blank or preform does not exceed 20%, and the average cooling rate of the at least one region from 700° C. to the Ms temperature is not less than 60° C./s.
  • the present invention proposes to achieve high strength through phase transformation strengthening, that is, after hot stamping deformation and die cooling, the component can obtain as much as possible Martensitic structure, or even full martensitic structure (total amount exceeds 97%). Based on the above considerations, in the low carbon design, how to make the steel plate obtain sufficient hardenability through the ratio of alloy elements is very important.
  • bainite has higher toughness than martensite.
  • this study found that in the case of low C content, the solid solution strengthening effect of carbon is weak, and a large amount of carbides will precipitate in the process of generating bainite, further reducing the solid solution strengthening effect of carbon in the matrix, thereby reducing the strength of the material. Therefore, the present invention intends to avoid bainite phase transformation as much as possible under low carbon conditions to ensure high strength.
  • the nucleation of bainite phase transformation is a diffusion-type phase transformation.
  • Si element can also promote the nucleation and growth of ferrite, expand the phase transformation temperature range of ferrite during cooling, and the formation of ferrite will also be detrimental to the high strength of the material.
  • the present invention under low-carbon design, on the one hand, controls the content of Cr and Si to inhibit bainite nucleation and reduce the promotion of ferrite nucleation respectively, and on the other hand, controls hardenability by 0.26Si+1.3Cr+Mn.
  • the combination of the two obtains a reasonable ratio of alloy elements such as Mn, Si and Cr, thereby ensuring as much martensitic structure as possible.
  • This is beneficial to the production control of hot stamping components on the one hand, and reduces the proportion of non-martensitic phases such as ferrite, bainite, and residual austenite on the other hand, thereby maximizing the yield strength and tensile strength of the component while obtaining sufficient plasticity and toughness.
  • the low Ac3 temperature makes it easy for the steel sheet to achieve complete austenitization under the heating conditions in the hot stamping method. Therefore, by controlling the Ac3 temperature, it is beneficial to further reduce the ferrite content in the final structure.
  • the appropriate Ms temperature makes the final martensitic matrix have both high strength and toughness.
  • increasing the Ms temperature is conducive to achieving martensite self-tempering and improving martensite toughness.
  • this study found that when the C content is low, the low-carbon martensitic matrix already has good toughness, but if the Ms is higher than 410°C, it will lead to excessive self-tempering, resulting in a decrease in the strength and toughness of the martensitic matrix.
  • the Ms temperature when the C content is low, it is preferred to control the Ms temperature within the range of 350°C to 410°C to ensure that the degree of martensite self-tempering can be effectively controlled so that the final martensitic matrix has both high strength and toughness.
  • the hot stamping steel sheet of the present invention has improved hardenability; and the hot stamping component of the present invention has a yield strength YS of 880-1100MPa and a tensile strength TS of not less than 1000MPa, an elongation of not less than 5%, a fracture strain ⁇ of not less than 0.6, and a bending energy absorption W/t 2 ⁇ 2.88 ⁇ 10 4 N/mm.
  • the high yield strength and tensile strength prevent the component from premature deformation and failure during the collision while ensuring the lightweight effect of the component, and the combination of high strength and excellent toughness and plasticity ensures the energy absorption effect of the component during the collision.
  • Figure 1 shows the typical microstructure of a hot stamped component with T1 composition after hot stamping.
  • Figure 2 shows the typical microstructure of a hot stamped component with T2 composition after hot stamping
  • Figure 3 Typical microstructure of hot stamped components with CT1 composition after hot stamping.
  • M represents martensite and B represents bainite.
  • the steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage, 0.053% ⁇ C ⁇ 0.10%, 0.05% ⁇ Si ⁇ 0.30%, 1.81% ⁇ Mn ⁇ 2.7%, 0.01% ⁇ Cr ⁇ 0.7%, 0.01% ⁇ Al ⁇ 0.5%, 0.0005% ⁇ B ⁇ 0.005%, 0.015% ⁇ Ti ⁇ 0.05%, 0 ⁇ Nb+V ⁇ 0.2%, 0.001% ⁇ P ⁇ 0.
  • the microstructure of the component comprises: less than 5% of bainite, less than 3% of austenite, less than 3% of ferrite and less than 0.2% of Nb-V-Ti microalloy carbide, and the remainder is martensite.
  • the steel plate matrix further comprises, in mass percentage, at least one of: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01%, and 0.0001% ⁇ REM ⁇ 0.01%.
  • the balance of the steel plate matrix is Fe and unavoidable impurities.
  • 0.03% ⁇ Ti ⁇ 0.046% is preferred.
  • C is the most commonly used alloying element in steel to improve strength. The higher the C content, the better the hardenability of the steel and the higher the strength. However, as the C content increases, the fracture strain of the steel plate decreases, resulting in a decrease in the toughness of the steel plate. At the same time, C also significantly affects the phase change characteristics of the steel, and a reduction in the C content will significantly increase the Ac3 temperature and Ms temperature of the steel. As the basic design of the present invention, the present invention requires that the C content be controlled not to exceed 0.10% so that the hot stamping formed components have the expected toughness. However, since C itself has the function of enhancing the work hardening ability of the material during deformation, the addition of a certain content of C is also very important. Based on the above considerations, the present invention requires that the C content be controlled within the range of 0.053 to 0.10%, preferably within the range of 0.055 to 0.09%.
  • Si 0.05-0.30%
  • Mn 1.81-2.70%
  • Cr 0.01-0.7%
  • the present invention needs to design the ratio of elements such as Si, Mn and Cr that improve hardenability and economy.
  • the present invention requires the addition of more than 0.05% Si.
  • more than 0.30% Si not only causes more serious surface oxidation and decarburization during the production process, thereby affecting the surface quality of the final product, but also because of its role in stabilizing ferrite, it is not conducive to obtaining as much martensitic structure as possible after hot stamping of the material, which is in turn not conducive to strength. Therefore, the present invention requires that the Si content is in the range of 0.05 to 0.30%, preferably in the range of 0.10 to 0.26%.
  • Mn improves the hardenability of steel, increases the stability of austenite, expands the austenitization phase area and reduces the Ac3 temperature.
  • Mn content is lower than 1.81%, it cannot offset the increase in Ac3 temperature and Ms temperature caused by the low C content in the present invention, resulting in the inability to achieve full austenitization under the heating conditions in the hot stamping method of the present invention, and more ferrite or bainite will be generated under the cooling conditions in the hot stamping method of the present invention, thereby reducing the strength and toughness of the component.
  • the Mn content is higher than 2.70%, segregation is prone to occur in steel during production, which has an adverse effect on the ductility and toughness of the steel. Therefore, the present invention requires that the Mn content be controlled within the range of 1.81 to 2.70%, preferably within the range of 2.01 to 2.50%.
  • Cr can not only improve the hardenability of steel, but also has a significant effect on the oxidation resistance of steel and the prevention of surface decarburization.
  • the present invention requires that when Si, Mn and Cr are within the above ranges, Mn+0.26Si+1.3Cr ⁇ 2.20%, preferably not less than 2.25%, and more preferably not less than 2.50%.
  • the present invention requires that the sum of Si+Cr is not higher than 0.50%, and more preferably, not higher than 0.47%.
  • Al is a strong deoxidizing element, so it is often added in an amount of not less than 0.01% as a deoxidizer during steel smelting. Excessive Al will increase the Ac3 temperature of the steel, resulting in the failure of full austenitization under the given hot stamping heating conditions of the present invention. At the same time, it will also cause the problem of increased resistance at the crystallizer mouth during continuous casting. Therefore, the Al content in the present invention does not exceed 0.5%.
  • the B element can be segregated at the austenite grain boundary, inhibiting the formation of ferrite and significantly improving the hardenability of the steel. Therefore, a certain amount of B is added to the components of the present invention. However, B cannot be added excessively, as too high a B content will lead to boron embrittlement. Therefore, the B content of the steel of the present invention is controlled at 0.0005-0.005%.
  • Ti and N have a strong binding force. Adding not less than 0.015% of Ti in hot stamping steel will form TiN with N in solid solution in the steel, avoiding the combination of N and B element and reducing the effect of B. At the same time, Ti forms carbonitrides, which are helpful for improving strength and grain refinement to improve toughness. However, it is not advisable to add too much Ti. Exceeding 0.05% will cause the toughness of the material to decrease. Therefore, the content of Ti is controlled at 0.015-0.05% in the present invention.
  • Nb and V in steel will form carbonitrides, which improve the performance of the material through precipitation strengthening and grain refinement. Therefore, the components described in the present invention can be added in appropriate amounts. However, excessive addition of the above elements will lead to an increase in material production costs. Therefore, the present invention controls the sum of the contents of Nb and V to be 0 to 0.2%, preferably 0.01 to 0.2%.
  • P is an unavoidable element.
  • P as a solid solution strengthening element can improve the strength of steel plates relatively cheaply.
  • the upper limit of the P content is not more than 0.100%, preferably not more than 0.050%.
  • the lower limit of the P content is not less than 0.001%, preferably not less than 0.004%.
  • the upper limit of the S content is not more than 0.100%, preferably not more than 0.015%.
  • the lower limit of the S content is not less than 0.0001%, preferably not less than 0.0005%, and more preferably not less than 0.001%.
  • the addition of W, Mo, Ni, Cu, and Co can improve the hardenability of steel, but the addition of these alloy elements will increase the cost of the alloy, so only an appropriate amount is added to the material.
  • Ni, Cu, and Co also have the benefit of improving the toughness of the material.
  • the content of each element is not less than 0.01%, the above-mentioned beneficial effect can be exhibited. Therefore, preferably, the lower limit of the content of the five elements is not less than 0.01%, respectively.
  • the upper limit of the content of the five elements is not more than 0.30%, respectively. In this case, while being able to improve toughness, it is ensured that the hardenability of the steel is less affected, and the machinability of the steel is guaranteed.
  • the content of Sn and Sb is not less than 0.005%, the wettability of the coating can be improved. Therefore, the lower limits of the contents of Sn and Sb are preferably not less than 0.005%, respectively. However, when more than 0.300% of Sn and/or more than 0.100% of Sb are contained, the toughness of the material will deteriorate. Therefore, the content of Sn is preferably not more than 0.300%, and the content of Sb is preferably not more than 0.100%.
  • REM Reare Earth Metal
  • Ca, Mg, and Zr can achieve the effect of refining inclusions by having a content of not less than 0.0001%, thereby improving the performance of the material. Therefore, the contents of Ca, Mg, Zr, and REM are preferably not less than 0.0001%, respectively.
  • the contents of Ca, Mg, Zr, and REM are preferably not more than 0.01%, respectively.
  • the sum of the contents of W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM elements is in the range of 0.0001% to 0.30%.
  • the steel plate matrix of the present invention may contain elements such as W, Mo, Ni, Cu, Co, Sn, Sb, Ca, Mg, Zr and REM, and the presence of these elements does not affect the solution of the technical problem of the present invention.
  • other components of the steel sheet matrix of the hot stamped component are not particularly limited.
  • elements such as As may be mixed in from scrap materials, they do not affect the characteristics of the hot stamped component if they are within a normal range.
  • the first example steel plate matrix of the hot stamping formed component of the present invention contains, by mass percentage, the following: 0.059% ⁇ C ⁇ 0.089%, 0.10% ⁇ Si ⁇ 0.30%, 1.88% ⁇ Mn ⁇ 2.47%, 0.08% ⁇ Cr ⁇ 0.31%, 0.015% ⁇ Al ⁇ 0.05%, 0.0015% ⁇ B ⁇ 0.0035%, 0.018% ⁇ Ti ⁇ 0.046%, 0 ⁇ Nb+V ⁇ 0.15%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, Fe ⁇ 95% and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy 2.20% ⁇ Mn+0.26Si+1.3Cr ⁇ 2.82%.
  • the first example steel plate matrix further comprises, by mass percentage, at least one of: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01%, and 0.0001% ⁇ REM ⁇ 0.01%. More preferably, 0.0001% ⁇ W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM ⁇ 0.30%.
  • the first example steel sheet matrix of the hot stamped component contains Fe and inevitable impurities as the balance in mass percent.
  • the second example steel plate matrix of the hot stamping formed component of the present invention comprises, by mass percentage, 0.059% ⁇ C ⁇ 0.089%, 0.10% ⁇ Si ⁇ 0.26%, 1.88% ⁇ Mn ⁇ 2.40%, 0.11% ⁇ Cr ⁇ 0.31%, 0.015% ⁇ Al ⁇ 0.05%, 0.0015% ⁇ B ⁇ 0.0035%, 0.018% ⁇ Ti ⁇ 0.046%, 0 ⁇ Nb+V ⁇ 0.15%, 0.001% ⁇ P ⁇ 0.100%, 0.0001% ⁇ S ⁇ 0.100%, Fe ⁇ 95% and unavoidable impurities, wherein the contents of Mn, Si and Cr satisfy 2.27% ⁇ Mn+0.26Si+1.3Cr ⁇ 2.60%.
  • the second example steel plate matrix further comprises, by mass percentage, at least one of: 0.01% ⁇ W ⁇ 0.30%, 0.01% ⁇ Mo ⁇ 0.30%, 0.01% ⁇ Ni ⁇ 0.30%, 0.01% ⁇ Cu ⁇ 0.30%, 0.01% ⁇ Co ⁇ 0.30%, 0.005% ⁇ Sn ⁇ 0.30%, 0.005% ⁇ Sb ⁇ 0.100%, 0.0001% ⁇ Ca ⁇ 0.01%, 0.0001% ⁇ Mg ⁇ 0.01%, 0.0001% ⁇ Zr ⁇ 0.01%, and 0.0001% ⁇ REM ⁇ 0.01%. More preferably, 0.0001% ⁇ W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM ⁇ 0.30%.
  • the second example steel sheet matrix of the hot stamped component contains Fe and inevitable impurities as the balance in mass percent.
  • the present invention enables the component to have sufficient hardenability, appropriate martensite transformation temperature and austenite transformation end temperature through alloy design.
  • the present invention hopes that the room temperature structure of the hot stamping formed component is as single martensite structure as possible.
  • the present invention makes the hot stamping formed component have a reasonable ratio of several economical alloying elements Mn, Si and Cr. Its content itself helps to reduce the production of non-martensite such as bainite and ferrite, and by satisfying Mn+0.26Si+1.3Cr ⁇ 2.20%, the component has sufficient hardenability.
  • the microstructure of the component of the present invention is composed of the following by area percentage: less than 5% bainite, less than 3% austenite, less than 3% ferrite, and the remainder is martensite.
  • the sum of the contents of bainite, ferrite and retained austenite is less than 5%. More preferably, the sum of the contents of bainite, ferrite and retained austenite does not exceed 3%.
  • the Ms temperature of the material is usually increased by alloy design, thereby improving the self-tempering effect of martensite, thereby improving the toughness of the hard and brittle martensitic structure of high C.
  • the C content of the steel of the present invention is relatively low, the interaction between carbon and dislocation in martensite is much weaker than that of high C content, so the toughness of martensite itself is better, and the effect of self-tempering on toughness improvement is limited.
  • the self-tempering degree of martensite is high, and martensite tempering decomposition softening and cementite precipitation are prone to occur, resulting in a significant reduction in the strength and toughness of the steel. Therefore, in the case of a low C content of the present invention, by using a high content of Mn and optimizing the ratio of alloying elements such as Si and Cr, the Ms temperature of the steel is controlled within the range of 350 to 410°C, preferably within the range of 360°C to 400°C, thereby controlling the self-tempering degree of the martensite matrix and achieving further optimization of the strength and toughness of the material.
  • the reduction of C content will also significantly increase the Ac3 temperature, resulting in the inability to obtain a complete austenite structure under the commonly used hot stamping heating conditions.
  • the present invention ensures that the Ac3 temperature is not higher than 900°C by using a high content of Mn and optimizing the ratio of alloying elements such as Si and Cr.
  • the present invention hopes to control the precipitation of microalloy carbides such as Nb, V, and Ti in the microstructure.
  • the precipitation of an appropriate amount of microalloy carbides will bring about grain refinement and precipitation strengthening effects, thereby improving the toughness of the component.
  • excessive carbide precipitation consumes carbon in martensite, thereby reducing the interaction between carbon and dislocations in martensite, which is not conducive to obtaining high strength of the component. Therefore, it is necessary to limit the content of Nb, V, and Ti to the aforementioned range.
  • the hot stamping formed component of the present invention may also be coated with a metal coating on its surface, and the coating may be an aluminum-based alloy coating or a zinc-based alloy coating.
  • the high strength of the material helps to improve the lightweight effect of hot stamping components.
  • a hot stamping component fails in collision, the component first undergoes local plastic deformation at the collision site, followed by bending and cracking in the deformation area. Therefore, the yield strength of the material becomes an important evaluation basis, which represents the critical strength of the steel material from elastic deformation to plastic deformation.
  • High yield strength can significantly delay the plastic deformation stage of the component, prevent the component from premature plastic deformation, and facilitate the repair of parts after collision.
  • the bending performance of the material defines the bending toughness of the steel plate. Good bending performance delays the time when the component fails during bending deformation and avoids early brittle fracture. Therefore, the component can absorb energy through continuous deformation, and the collision energy absorption performance will be greatly improved.
  • the commonly used method for testing material toughness is the static three-point bending test (i.e., VDA bending test, VDA 238-100 standard).
  • VDA bending test VDA 238-100 standard.
  • the test can obtain the maximum bending angle, the peak force corresponding to the maximum bending angle, the bending energy absorption, and the fracture strain, thereby reflecting the material's ability to resist bending deformation failure.
  • the room temperature tensile test (GB/T 228.1 standard) is a common method for measuring material strength. The yield strength and tensile strength reflect the material's ability to resist tensile deformation failure.
  • the bending moment in the deformation area is proportional to the square of the steel plate thickness
  • the bending energy absorption W represents the sum of the component's ability to generate bending moment during the entire VDA bending deformation process, and is also proportional to the square of the steel plate thickness. Therefore, for specimens of different thicknesses, in order to eliminate the energy difference caused by the plate thickness, W/t 2 is used to characterize the unit bending energy absorption performance.
  • the component of the present invention has a yield strength YS of 880-1100 MPa and a tensile strength TS of not less than 1000 MPa, an elongation of not less than 5%, and a fracture strain ⁇ of not less than 0.6.
  • the component of the present invention has a yield strength of 930-1050 MPa and a tensile strength of not less than 1060 MPa, an elongation of not less than 7%, and a fracture strain of not less than 0.7.
  • the bending energy absorption W of the static three-point bending of the component of the present invention is not less than 50000 mm ⁇ N, and there is the following relationship between W and the plate thickness t: W/t 2 ⁇ 2.88 ⁇ 10 4 N/mm.
  • the present invention provides a hot stamping method for producing the hot stamping formed component, which comprises:
  • step B) is completed in a heating furnace, and the maximum temperature in the furnace is in the range of 900-980°C to achieve complete austenitization.
  • hot stamping step transferring the heated blank or preform to a press for hot deformation and cooling to below 300°C in a die to obtain the hot stamping component, wherein the average cooling rate of the hot stamping component from 700°C to the Ms temperature is not less than 40°C/s in the area where no hot deformation occurs, and the hot stamping component is cooled at a higher average cooling rate in the area where hot deformation occurs.
  • the Ms temperature of the hot stamping component satisfies: 350°C ⁇ Ms ⁇ 410°C.
  • a high average cooling rate can minimize the formation of non-martensite phases (such as bainite, ferrite, and retained austenite), thereby ensuring the strength and toughness of the hot stamping formed components.
  • non-martensite phases such as bainite, ferrite, and retained austenite
  • the deformation amount in at least one thermal deformation area of the component does not exceed 10% and the average cooling rate of the thermal deformation area is not less than 50°C/s. More preferably, in step C, the thermal deformation amount of at least one area of the blank or preform does not exceed 20%, and the average cooling rate of the at least one area from 700°C to the Ms temperature is not less than 60°C/s.
  • the steel is prepared into experimental steel plates by the following process, namely:
  • Hot rolling The steel billet is heated to 1200°C and kept at this temperature for 2h, then hot rolled at 800°C ⁇ 1200°C, and coiled at below 700°C to form hot rolled steel coil, and the hot rolled coil is pickled to remove the oxide scale generated during the hot rolling process;
  • Cold rolling cold rolling the pickled hot-rolled coil with a cold rolling reduction of 30% to 70% to obtain a cold-rolled steel coil with a thickness of 1.4 to 1.9 mm;
  • Hot-dip coating hot-dip aluminum alloy coating is applied to cold-rolled sheets of T1, T2, T4, T5, T7, T8, T9, T10, T11 and CT1-CT3 to obtain the final test steel sheets, with a single-side coating weight of 20-75 g/m 2 ; and hot-dip zinc alloy coating is applied to cold-rolled sheets of T6 to obtain the final test steel sheets, with a single-side coating weight of 40-80 g/m 2 ;
  • T1 to T11 are exemplary steels within the composition range of the present invention and satisfying the ratio relationship of Mn, Si and Cr.
  • CT1 to CT3 are exemplary steels with the composition range of the present invention and satisfying the ratio relationship of Mn, Si and Cr. It is a comparative steel outside the present invention.
  • the C content of T1 to T11 is in the range of 0.059 to 0.089%, the value of Mn+0.26Si+1.3Cr is not less than 2.20%, and the value of Si+Cr is not greater than 0.47%.
  • the calculated Ac3 temperature of T1 to T11 is not higher than 900°C, and the calculated Ms temperature is in the range of 366 to 396°C.
  • the Si content of CT1 is high and the Mn content is low, so that the calculated Mn+0.26Si+1.3Cr is lower than 2.20, the hardenability is poor, and high Si promotes the formation of ferrite, which is easy to generate non-martensite phase, which is not conducive to improving strength.
  • the calculated Ms temperature is also high, reaching 417°C, so that the degree of martensite self-tempering cannot be effectively controlled, and it is not conducive to improving strength.
  • the C and Mn contents of CT2 are high, resulting in a low Ms temperature, and martensite self-tempering cannot be effectively achieved to improve martensite toughness.
  • the composition range of CT3 is within the composition range of the present invention, but the calculated Mn+0.26Si+1.3Cr is lower than 2.20, indicating poor hardenability.
  • the Si content is within the scope of the present invention, the hardenability is not enough to inhibit the promotion of Si to ferrite, so that a small amount of ferrite will still be generated, which is not conducive to improving strength.
  • the calculated Ac3 temperature is also high, reaching 906°C, which is not conducive to obtaining complete austenite in the hot stamping heating process.
  • more Si and Cr were added to CT3, and the value of Si+Cr reached 0.52%.
  • the value of Si+Cr should not exceed 0.50%.
  • critical cooling rate The critical cooling rate (hereinafter referred to as "critical cooling rate") for complete martensitic transformation of components T1, T2, T4, T5 and CT1 was tested by DIL805 phase transformer. Under normal circumstances, the greater the critical cooling rate, the higher the cooling requirement for complete martensitic transformation of the material, that is, the worse the hardenability of the material.
  • samples were heated to 930°C for insulation. Without deformation and 10% thermal deformation, they were cooled to below 200°C at a cooling rate of 20-70°C/s to obtain the expansion curve in the cooling process (from 700°C to 400°C). It is determined whether there is a non-martensitic phase transformation on this section of the expansion curve and the inflection point temperature is recorded.
  • the minimum cooling rate without non-martensitic phase transformation on the curve is taken as the critical cooling rate of the material.
  • the recorded results are shown in Table 2.
  • the critical cooling rate of the martensitic phase transformation is determined based on a deformation of 10%.
  • the Mn+0.26Si+1.3Cr calculated from T1 and T2 are both 2.20, and the hardenability of T1 and T2 is poor compared to other example steels, such as T4 (whose Mn+0.26Si+1.3Cr is 2.27) and T5 (whose Mn+0.26Si+1.3Cr is 2.53). Therefore, if T1 and T2 can meet the cooling requirements, the other components will also meet the requirements, so the test results of T1 and T2 are used as the benchmark for the critical cooling rate.
  • the critical cooling rate of T5 is below 20°C/s
  • the critical cooling rate of T4 is around 25°C/s
  • the critical cooling rates of T1 and T2 are around 30°C/s.
  • the critical cooling rate of the commonly used 22MnB5 material is about 27°C/s. This shows that through the optimized design of Mn, Si and Cr, the value of Mn+0.26Si+1.3Cr is not less than 2.20%. Even if the C content is low, the steel of the present invention still has good hardenability. Therefore, without changing the existing production conditions of hot stamping, it can ensure that a full martensitic structure is obtained during the cooling process.
  • CT1 since the value of Mn+0.26Si+1.3Cr is about 1.70%, in the absence of thermal deformation, even when the cooling rate reaches 40°C/s, CT1 still has significant non-martensitic phase transformation during the cooling process, and the hardenability is poor. Therefore, under the existing hot stamping forming conditions, CT1 may generate more ferrite and bainite, which is not conducive to production control and the strength guarantee of the final parts.
  • the critical cooling rates of T5 and T4 are in the range of 25-30°C/s and 40-50°C/s, respectively, while the critical cooling rate of T2 is in the range of 50-60°C/s. That is, the lower the value of Mn+0.26Si+1.3Cr, the higher the critical cooling rate required.
  • the values of Mn+0.26Si+1.3Cr for T1 and T2 are the same, but T1 has a lower critical cooling rate than T2, in the range of 40-50°C/s°C. This is because more Cr, which is beneficial to inhibit the formation of bainite, and less Si, which promotes the formation of ferrite, are added to T1, which helps to reduce the critical cooling rate. Therefore, when the hardenability is close without hot deformation, the combined addition of higher Cr content and lower Si content is beneficial to obtain a complete martensitic structure of the component under hot deformation conditions.
  • the case of hot deformation is compared with the case without hot deformation.
  • a higher critical cooling rate will be required.
  • the component can obtain a uniform martensitic structure.
  • the average cooling rate can be further reduced to above 50°C/s, so that the component obtains a uniform martensitic structure.
  • the steel of the present invention improves the hardenability of the steel by reasonably matching Mn, Si and Cr on the basis of ensuring performance, inhibits the formation of non-martensitic phases, and reduces the requirements for production control.
  • the hot stamping hat-shaped parts were generated by hot stamping of T1 ⁇ T11 and CT1-CT3 steel plate samples of different thicknesses: the sample with a size of 230*550mm was heated to 910°C and kept warm to completely austenitize, and then the sample was transferred to the press for hot stamping, and the transfer time was 10 ⁇ 12s. After that, the hat-shaped component was taken out after hot stamping in the press and cooling for 8 ⁇ 10s to below 250°C. The average cooling rate of the sample during the cooling process can reach more than 50°C/s.
  • the bending fracture strain test method is as follows: (1) Determine the VDA bending angle ⁇ peak of the component specimen by a static three-point bending test; (2) Based on the experimental results, select at least three groups of interruption bending angles ⁇ L (i.e., the bending angle of the component specimen under load) for interruption bending tests to ensure that ⁇ L ⁇ 50% ⁇ peak ; (3) Stop loading when the component specimen is bent to ⁇ L , and measure the bending angle ⁇ UL of the component specimen under unloading; (4) Place the unloaded component specimen under an optical microscope, and measure the inner and outer surface radii R i and R o of the most severe deformation zone; (5) Calculate the equivalent (plastic) strain ⁇ of the outer surface of the most severe deformation zone of the component specimen under unloading conditions of different ⁇ UL according to equation (1), that is, the equivalent strain of the outer surface of the most severe deformation zone of the component specimen when the component specimen is bent to ⁇ L , thereby establishing the ⁇ - ⁇ L relationship;
  • Figure 1 shows the typical microstructure of a T1 component sample of the present invention, in which the martensite phase is dominant and the total amount of non-martensite phase is less than 3% (area percentage).
  • the typical microstructure of the T1 component sample is applicable to T3-T11 component samples.
  • Figure 2 is a typical microstructure of the T2 component sample of the present invention, in which a small amount of bainite exists in its microstructure, which is less than 5% by area percentage.
  • Figure 3 is a typical microstructure of the CT1 component sample, in which a large amount of bainite exists in its microstructure, which is more than 10% by area percentage.
  • the yield strength of the T1-T11 component samples of the present invention is in the range of 880-1041 MPa
  • the tensile strength is in the range of 1054-1186 MPa
  • the elongation is in the range of 7-8%
  • the maximum bending angle reaches 88-99.7°
  • the fracture strain is above 0.7
  • the W/t 2 value is not less than 2.9 ⁇ 10 4 N/mm, showing an excellent match between strength and toughness.
  • the yield strength, tensile strength, peak force and W/t 2 value have all decreased significantly, which are 816MPa, 961MPa, 6568N and 2.84 ⁇ 10 4 N/mm respectively, which do not meet the use requirements.
  • the structure of the CTI component specimen contains more than 10% bainite and a small amount of ferrite.
  • bainite can improve the toughness of the material. Therefore, the maximum bending angle of the CT1 component specimen reaches 104° and the fracture strain is 0.90. However, the presence of more bainite and a small amount of ferrite is not conducive to the strength of the material.
  • the high Ms temperature causes excessive self-tempering of martensite, which also leads to a decrease in strength.
  • the low Cr content (0.04%) is not enough to inhibit the formation of bainite and the high Si content (0.42%) promotes the formation of ferrite, which further leads to the formation of more bainite and ferrite in the microstructure under the same hot stamping forming conditions.
  • the toughness of the CT2 component sample is lower than that of all other component samples, with a maximum bending angle of only 79.3° and a fracture strain of 0.59. The reason is that the Mn content is higher than 2.70%, which will cause segregation, thereby reducing the ductility and toughness of the steel.
  • the toughness of the CT3 component sample is good, but the yield strength is reduced by about 7%, failing to reach 880MPa. This is because although the components of both are within the scope of the present invention, the Mn, Si and Cr ratio of the CT3 component sample fails to meet Mn+0.26Si+1.3Cr ⁇ 2.20%, resulting in poor hardenability and the generation of relatively more bainite (area percentage of about 5-10%), so its yield strength is low.
  • each alloying element in order to obtain the expected high strength and high toughness, such as a tensile strength of not less than 880 MPa, a maximum bending angle of not less than 80°, a fracture strain of not less than 0.6 and a W/t 2 value of not less than 2.88 ⁇ 10 4 N/mm, each alloying element must not only be within the scope of the present invention, but also Mn, Si and Cr must satisfy Mn+0.26Si+1.3Cr ⁇ 2.20%, and both are indispensable.
  • Mn+0.26Si+1.3Cr 2.25% or more
  • a yield strength of not less than 930MPa can be obtained.
  • the Mn+0.26Si+1.3Cr of the T5 component sample is 2.53%, and the yield strength reaches 954MPa. Therefore, when Mn+0.26Si+1.3Cr is 2.50% or more, a yield strength of not less than 950 MPa can be obtained.
  • the toughness of the T1 component sample is close to that of the T2 component sample, but the yield strength and tensile strength of the T2 component sample are significantly lower than those of the T1 component sample. This is because although the two have similar C contents and the same value of Mn+0.26Si+1.3Cr, the Cr content of the T2 component sample is lower than that of the T1 component sample, so it is not enough to suppress the generation of bainite during the cooling process, so that the obtained structure contains a small amount of bainite (see Figure 2). The presence of bainite reduces the yield strength and tensile strength of the T2 component sample.
  • CT1 also has a large amount of bainite due to the low Cr content, so the yield strength and tensile strength are significantly reduced. Therefore, while ensuring hardenability and toughness, the strength can be further improved by increasing the Cr content.
  • T3 and T6 component samples are uncoated and galvanized samples, respectively, and all the properties of the two meet the use requirements. It can be seen that whether it is galvanized, aluminum-plated or uncoated, the obtained hot stamping formed components can meet the expected performance requirements.
  • the experimental results of T1 to T11 also show that the sample components with different thicknesses of the present invention can meet the expected performance requirements.
  • the hot stamping formed components obtained can meet the expected performance requirements regardless of whether the zinc coating layer, the aluminum coating layer or the no coating layer.
  • the requirements of high strength, high toughness and high energy absorption capacity are met, which improves the safety of use.
  • the performance can be further improved by further controlling the range of alloy elements.

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Abstract

一种低碳的高韧性热冲压成形构件及钢板。热冲压成形构件的钢板基体以质量百分数计包含0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，Mn+0.26Si+1.3Cr≥2.20％；以面积百分数计钢板基体的微观组织包括小于5％贝氏体、小于3％奥氏体、小于3％铁素体及小于0.2％Nb-V-Ti碳化物，剩余为马氏体。

Description

一种低碳的高韧性热冲压成形构件及钢板 技术领域
本发明涉及一种低碳的高韧性热冲压成形构件及钢板。
背景技术
作为汽车钢中强度级别最高的材料品类，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。室温拉伸试验(GB/T 228.1标准)是一种常见的测量材料强度的方法，所获得的屈服强度和抗拉强度反映了材料抵抗拉伸变形失效的能力。同时，检测材料韧性比较常用的方法是静态三点弯曲试验(即VDA弯曲试验，VDA 238-100标准)。通过试验可以有效地测量材料的断裂应变，从而反映出材料抵抗弯曲变形失效的能力。
为了满足汽车安全性，通常要求汽车的车身的特定区域在碰撞时具有充足的吸能效果，尤其是在侧面碰撞时。在这些特定区域(如B柱的下端)中，要求在汽车碰撞过程中，通过零件变形吸收大量的冲击能量且零件不发生脆性断裂。同时，该零件还需要兼具高的屈服强度和加工硬化能力，以防止零件过早发生变形失效。此外，高强度的零件能够更好地实现汽车的轻量化，以满足节能环保的需求。目前，在前述特定区域中的零件普遍使用热冲压钢22MnB5制成。尽管这种钢板拥有1400MPa以上的抗拉强度，但其韧性偏低，难以满足日益趋严的碰撞法规和变形吸能的要求，存在安全隐患。
CN107810281A提供了一种用于压制硬化的钢和由这样的钢制造的压制硬化的部件。所述压制硬化的钢部件主要用于机动车辆中的用于防侵入或能量吸收功能的结构元件。通过使钢中的碳含量满足0.062％≤C≤0.095％来获得令人满意的拉伸强度、弯曲性和焊缝韧性，并且通过使钢中的淬透性元素C、Mn、Si、Cr的含量满足1.5％≤(C+Mn+Si+Cr)≤2.7％来保证热冲压部件的显微组织中马氏体和提高延性的贝氏体的含量配比。该专利教导利用较多的贝氏体来提高部件的韧性和可弯曲性。因此，该专利技术特别要求Cr的含量不高于0.1％，因为Cr的添加不利于提高延性的贝氏体的形成。然而，一方面，现有热冲压生产过程难以获得所要求比例的贝氏体和马氏体结合的混合组织；另一方面，较多贝氏体的存在对于部件的强度，特别是屈服强度的提高是不利的。低的屈服强度将会造成部件在发生碰撞时过早变形失效，同样降低了车辆的安全性。
为了提高韧性，CN104769138A发现，在涂覆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热成形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。通过这种方法，1.8mm厚的镀锌22MnB5热成形钢板的VDA弯曲角可以达到55°以上。但是上述特定厚度的脱碳区很大程度上依赖于涂覆镀层前的退火条件，从而提高了生产难度。另外，该专利忽略了脱碳层对材料的强度、表面硬度、弯曲时的峰值载荷以及后续的涂镀质量的不利影响。
CN1717499A提供了一种用于冷加工成形的高强度钢板及制造方法。为保证钢板的优异加工性和高强度，控制该钢板的成分为C：0.05～0.15％，Si：0.3～2.0％，Mn：1.0～2.8％，Al：0.005～0.5％，N：0.006％以下，其余为Fe和不可避免杂质，并且Mn/C≥12且Si/C≥4。C、Mn及Si之间的比例设计用以延缓涂镀后的合金化处理过程中珠光体和渗碳体相变。该专利的钢板的显微组织以铁素体为主并含有3～20％的马氏体和残余奥氏体，其抗拉强度为500～900MPa，延伸率为27～38％，兼顾了高强度和高加工性。然而，由于热冲压成形过程通常涉及在870～930℃保温并且之后模具冷却，所以在经历热冲压成形过程时无法保证该高强度钢板能够获得该专利提到的理想的显微组织，从而无法获得高强度和高加工性。此外，针对合金化过程(温度一般为450～520℃)设计的Mn/C≥12和Si/C≥4在热冲压成形过程(不包含合金化过程)中也无法实现其预期目的。因此，CN1717499A的高强度钢板并不适用于制备具备高强度和高韧性的热冲压成形构件。
鉴于上述，为了提高汽车安全性，仍存在开发一种热冲压成形用钢，具备高强度和高韧性的热冲压成形构件及热冲压成形方法的需求。
发明内容
本发明是鉴于现有技术中存在的上述问题做出的。
本发明的目的之一在于在低的C含量的前提下，提供一种具有改善的淬透性的热冲压成形用钢板。本发明的热冲压成形用钢板的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
优选地，所述钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％、0.0001％≤REM≤0.01％中的至少一者。
进一步优选地，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
优选地，本发明的热冲压成形用钢板的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，余量为Fe及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
优选地，所述钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，及下述中的至少一者：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％和0.0001％≤REM≤0.01％，且0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％，余量为Fe及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
本发明在低的C含量的前提下，通过Mn、Cr、Si等元素合理配比获得了改善的淬透性。
优选地，Mn+0.26Si+1.3Cr≥2.25％，进一步Mn+0.26Si+1.3Cr≥2.50％。
优选地，0.11％≤Cr≤0.5％；更优选地，0.12％≤Cr≤0.4％。
优选地，2.01％≤Mn≤2.50％。
优选地，0.01≤Nb+V≤0.2％。
优选地，0.055％≤C≤0.09％且0.10％≤Si≤0.26％。
优选地，Si+Cr≤0.50％，更优选地Si+Cr≤0.47％，以进一步改善镀层钢板的涂镀表面质量，同时，使得热冲压后获得更均匀的组织和更优异的钢板性能。
在本文中，所述钢板的马氏体相变开始温度Ms通过以质量百分数计的合金元素计算为Ms＝520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu。优选地，满足：350℃≤Ms≤410℃，进一步优选地，360℃≤Ms≤400℃。Ms的计算公式由A.V.Sverdlin和A.R.Ness的《Steel Heat Treatment Handbook》(G.E.Totten和M.A.H.Howes,Ed.,Marcel Dekker Inc.，New York，1997)中获得。
在本文中，所述钢板的奥氏体转变结束温度Ac3通过以质量百分数计的合金元素计算为Ac3＝912-250C-16Mn+48Si-2Cr-16Ni+95V+96Ti+210Al-10Cu。优选地，满足：Ac3≤900℃。Ac3的计算公式由曾其英的《钢中合金元素对相变点Ac1、Ac3影响的定量关系探讨》 (理化检验-物理分册，1985，18(5):47-49)中获得。
所述热冲压成形用钢板可具有金属镀层。优选地，所述金属镀层可以是基于铝的合金镀层，也可以基于锌的合金镀层。
本发明的另一目的在于提供一种具有高强度和高韧性的热冲压成形构件。本发明的热冲压成形构件的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％；其中，以面积百分数计，所述钢板基体的微观组织包括：小于5％的贝氏体、小于3％的奥氏体、小于3％的铁素体以及小于0.2％的Nb-V-Ti微合金碳化物，剩余部分为马氏体。
优选地，所述热冲压成形构件的钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％、0.0001％≤REM≤0.01％中的至少一者。
进一步优选地，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
优选地，本发明的热冲压成形构件的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，余量为Fe及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
优选地，本发明的热冲压成形构件的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，及下述中的至少一者：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％和0.0001％≤REM≤0.01％，其中，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％，余量为Fe及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
与现有技术相比，在低的C含量的前提下，本发明的热冲压成形构件通过合理配比Mn、Cr、Si等多种合金元素被优化，从而能够在保证高韧性的情况下实现高的强度，满足了车身的特定区域对碰撞吸能以及轻量化的需求。
优选地，贝氏体、铁素体及残余奥氏体的总和不超过3％。
所述构件可具有金属镀层。优选地，所述金属镀层可以是基于铝的合金镀层，也可以基于锌的合金镀层。
所述构件具有880～1100MPa的屈服强度YS和不低于1000MPa的抗拉强度TS，延伸率不低于5％，且断裂应变ε不小于0.6。
优选地，所述构件具有930～1050MPa的屈服强度和不低于1060MPa的抗拉强度，延伸率不低于7％，且断裂应变不小于0.7。
优选地，所述构件的静态三点弯的折弯能量吸收W与板厚t满足：W/t2≥2.88×104N/mm。
本发明的再一目的在于一种用于生产上述热冲压成形构件的热冲压方法，其包括：
A)提供上述热冲压成形用钢板的坯件或由所述热冲压成形用钢板预成形得到的预制件；
B)完全奥氏体化处理：将步骤A中的坯件或预制件加热至900℃以上并保温；
C)热冲压成形处理：在完成步骤B之后，将经加热的坯件或预制件转移到压机中进行热变形并在模具内冷却至300℃以下以获得热冲压成形构件，其中，热冲压成形构件在未发生热变形的区域中从700℃冷却至Ms温度的平均冷速不低于40℃/s，并且在发生热变形的区域中以更高的平均冷速进行冷却。
优选地，热冲压成形构件的Ms温度满足：350℃≤Ms≤410℃。优选地，步骤B在加热炉内完成，炉内的最高温度在900～980℃的范围内，并且坯件或预制件在炉内时间(分钟，包括加热和保温的时间)满足：坯件或预制件的厚度t(mm)～(t×3+10)。
更优选地，在步骤C中，坯件或预制件的至少一个区域的热变形量不超过10％，且该至少一个区域在从700℃冷却至Ms温度的平均冷速不低于50℃/s。
更优选地，在步骤C中，坯件或预制件的至少一个区域的热变形量不超过20％，且该至少一个区域在从700℃冷却至Ms温度的平均冷速不低于60℃/s。
众所周知，作为一种钢中常见的淬硬性元素，C的添加将显著提高钢材的强度，但不利于韧性。因此，为获得充足的韧性，降低钢材中的碳含量是最有效的方案。然而，低碳设计使得钢材的强度降低，不利于车辆轻量化的需求。为了兼顾钢材的强度和韧性，本发明提出通过相变强化来实现高强度，即经过热冲压变形和模具冷却后，构件能够获得尽可能多的 马氏体组织，甚至全马氏体组织(总量超过97％)。基于上述考虑，在低碳设计下，如何通过合金元素的配比使得钢板获得足够的淬透性至关重要。
另外，公知常识认为，贝氏体拥有比马氏体更高的韧性。但本研究发现，在低C含量的情况下，碳的固溶强化效果弱，而生成贝氏体的过程中大量碳化物将析出，进一步降低碳在基体中的固溶强化效果，从而使得材料强度降低。因此，本发明意欲在低碳情况下，尽可能避免发生贝氏体相变，以保证高强度。同时，本发明注意到贝氏体相变的形核为扩散型相变，Cr作为强碳化物元素可降低碳的扩散能力，故而能够抑制贝氏体形核，这跟其他单纯增加淬透性的合金元素(如Mn，热力学上的奥氏体稳定元素)的作用是不同的。此外，作为钢中的固溶强化元素，Si元素也可以促进铁素体的形核和长大，扩大了冷却过程中铁素体的相变温度区间，而铁素体的形成也将不利于材料的高强度。为此，本发明在低碳设计下，一方面通过控制Cr和Si的含量来分别抑制贝氏体形核和减少对铁素体形核的促进作用，另一方面通过0.26Si+1.3Cr+Mn控制淬透性，二者结合获得Mn、Si和Cr等合金元素的合理配比，从而保证尽可能多的马氏体组织。这一方面利于热冲压成形构件的生产管控，另一方面降低了铁素体、贝氏体、残余奥氏体等非马氏体相的比例，从而在获得充足的塑性和韧性的情况下，最大程度提高构件的屈服强度和抗拉强度。
再者，低的Ac3温度使得钢板在热冲压方法中的加热条件下能容易实现完全奥氏体化。因此通过控制Ac3温度，有利于进一步减少最终组织中铁素体的含量。
此外，适当的Ms温度使得最终的马氏体基体同时具有高的强度和韧性。虽然公知常识认为，提高Ms温度有利于实现马氏体自回火且改善马氏体韧性。但本研究发现，在C含量较低的情况下，低碳的马氏体基体已经具有良好的韧性，但是如Ms高于410℃，会导致过度的自回火，使得马氏体基体强度和韧性均下降。因此，本发明在C含量较低的情况下，优选将Ms温度控制在350℃～410℃范围内，以保证能够有效控制马氏体自回火的程度，使得最终的马氏体基体同时具有高的强度和韧性。
与现有技术相比，本发明的热冲压成形用钢板具有改善的淬透性；并且本发明的热冲压成形用构件具有880～1100MPa的屈服强度YS和不低于1000MPa的抗拉强度TS，延伸率不低于5％，断裂应变ε不小于0.6、且折弯能量吸收W/t2≥2.88×104N/mm。高的屈服强度和抗拉强度在保证构件的轻量化效果的同时阻止了构件在碰撞过程中发生过早变形失效，并且高强度与优异的韧塑性的结合保证了构件在碰撞过程中的能量吸收效果。
附图说明
图1经热冲压成形后，T1成分的热冲压成形构件的典型显微组织；
图2经热冲压成形后，T2成分的热冲压成形构件的典型显微组织；
图3经热冲压成形后，CT1成分的热冲压成形构件的典型显微组织。
各图中，M代表马氏体，B代表贝氏体。
具体实施方式
下面结合示例性实施例来更详细地描述本发明。本文中关于化学元素含量的描述(％)都是指重量百分数。除非特别指出，否则各优选方案可以按照需要自由组合。除非明确说明，否则所有范围均包括端值。本领域技术人员将理解，实施例中记载的数据和各种参数仅是示例性的，并不构成对本发明的限制。
本发明的热冲压成形构件的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％，其中，以面积百分数计，所述构件的微观组织包括：小于5％的贝氏体、小于3％的奥氏体、小于3％的铁素体以及小于0.2％的Nb-V-Ti微合金碳化物，剩余部分为马氏体。
优选地，所述钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％、0.0001％≤REM≤0.01％中的至少一者。
进一步优选地，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
优选地，所述钢板基体的余量为Fe及不可避免的杂质。
优选地，Mn+0.26Si+1.3Cr≥2.25％，进一步Mn+0.26Si+1.3Cr≥2.50％。优选地，0.11％≤Cr≤0.5％，更优选地，0.12％≤Cr≤0.4％。优选地，0.01≤Nb+V≤0.2％。优选地，0.055％≤C≤0.09％；0.10％≤Si≤0.26％。优选地，Si+Cr≤0.50％。优选地，1.90％≤Mn≤2.50％，进一步2.01％≤Mn≤2.50％。优选地，0.03％≤Ti≤0.046％。
优选地，所述构件的马氏体相变开始温度Ms通过以质量百分数计的合金元素计算为Ms＝520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu，且满足：350℃≤Ms≤410℃。所述构件的奥氏体转变结束温度Ac3通过以质量百分数计的合金元素计算为Ac3＝912-250C-16Mn+48Si-2Cr-16Ni+95V+96Ti+210Al-10Cu，且满足：Ac3≤900℃。
本发明所述热冲压成形构件的化学成分详细描述如下：
C：0.053～0.10％
C是钢中最常用的提高强度的合金元素。C含量越高，钢的淬硬性越好，强度越高。但是，随着C含量的提高，钢板的断裂应变降低，导致钢板的韧性降低。同时，C也显著影响钢的相变特性，C含量的降低将明显提高钢的Ac3温度和Ms温度。作为本发明的基础设计，本发明要求C含量控制不超过0.10％以使热冲压成形构件具有预期的韧性。然而，由于C本身具有提升材料变形过程中的加工硬化能力的作用，因此，一定含量C的添加也是非常重要的。基于上述考虑，本发明要求C含量控制在0.053～0.10％的范围内，优选地在0.055～0.09％范围内。
Si：0.05～0.30％，Mn：1.81～2.70％，Cr：0.01～0.7％且Mn+0.26Si+1.3Cr≥2.20％在低碳设计下，为保证热冲压过程中能够有尽可能多的马氏体生成，从而获得高强度，本发明需要对Si、Mn及Cr等提高淬透性且经济性的元素的配比进行设计。
Si固溶在基体中，起到提高强度的作用，也能够在一定程度上提高钢材的淬透性。因此，本发明需要添加0.05％以上的Si。然而，超过0.30％的Si不仅使得生产过程中会出现较严重的表面氧化和脱碳，从而影响最终产品的表面质量，而且还会由于其稳定铁素体的作用，不利于材料热冲压后获得尽可能多的马氏体组织，进而对强度不利。因此，本发明要求Si含量在0.05～0.30％范围内，优选地在0.10～0.26％范围内。
Mn提高了钢材淬透性，增加了奥氏体的稳定性，扩大奥氏体化相区并降低Ac3温度。当Mn含量低于1.81％时，不能抵消本发明中低的C含量导致的Ac3温度和Ms温度升高，从而导致在本发明的热冲压方法中的加热条件下不能实现全奥氏体化，且在本发明的热冲压方法中的冷却条件下会生成较多的铁素体或贝氏体，从而降低构件的强度和韧性。当Mn含量高于2.70％时，钢材在生产时容易出现偏析，这对钢材的延性和韧性均存在不利的影响。因此，本发明要求Mn含量控制在1.81～2.70％范围内，优选地2.01～2.50％范围内。
Cr不仅能够提高钢材淬透性，而且对钢材的抗氧化和防止表面脱碳具有显著效果。此外，发明人发现，Cr为强碳化物元素，具有降低碳扩散能力。由于贝氏体相变的形核机制为扩散相变，因此，添加不少于0.01％的Cr有助于抑制贝氏体相变，这是额外于其他增加淬透性的合金元素(如Mn)的作用。然而，超过0.7％的Cr在生产过程中会导致热轧卷形成较严重的表面氧化，影响热卷酸洗的表面质量。因此，本发明的Cr含量控制在0.01～0.7％范围内，优选地在0.11～0.5％范围内，更优选地，在0.12～0.40％范围内。
在Si、Mn及Cr满足上述范围的同时，当Mn+0.26Si+1.3Cr＜2.20％时，由于淬透性不足而导致强度无法满足要求。因此，本发明要求在Si、Mn及Cr在上述范围内的同时，满足Mn+0.26Si+1.3Cr≥2.20％，优选地不小于2.25％，且更优选地不小于2.50％。
此外，一方面，由于Si和Cr均是较Fe元素更易氧化的合金元素，大量添加将造成热卷表面严重氧化，降低酸洗效果，且影响涂镀板的涂镀质量及其热冲压加热过程中的Fe和Al元素在镀层和基材界面处的相互扩散效果。另一方面，Si和Cr的合适配比有利于抑制热冲压后铁素体和贝氏体等非马氏体的相组织的生成，从而获得尽可能多的马氏体组织和更优异的综合性能。因此，优选地，本发明要求Si+Cr之和不高于0.50％，更优选地，不高于0.47％。
Al：0.01～0.5％
Al是强脱氧元素，因此，常添加不小于0.01％的量用作钢材熔炼时的脱氧剂。过量的Al会提高钢材的Ac3温度，导致在本发明给定的热冲压加热条件下不能发生全奥氏体化，同时，还会造成连铸时结晶器口阻力增大的问题。因此，本发明中的Al含量不超过0.5％。
B：0.0005～0.005％
在热冲压过程中，B元素能够偏聚在奥氏体晶界处，抑制铁素体的生成，显著提高钢的淬透性，因此，本发明的构件中均添加了一定量的B。但B也不能过量添加，B含量过高会导致硼脆，因此，本发明钢材B含量控制在0.0005～0.005％。
Ti：0.015～0.05％
Ti与N具有很强的结合力，热冲压钢中添加不小于0.015％的Ti，会与钢中固溶态的N形成TiN，避免了N与B元素的结合，降低B的作用。同时，Ti形成碳氮化物，对于强度的提高以及晶粒细化改善韧性具有帮助。但Ti不宜添加过多，超过0.05％会导致材料韧性下降。因此，本发明中控制Ti的含量在0.015～0.05％。
Nb+V：0～0.2％
Nb、V在钢中的添加，将形成为碳氮化物，通过析出强化和晶粒细化等作用改善材料的性能，因此，本发明中所述的构件可适量添加。然而，上述元素的过量添加，则会导致材料生产成本上升。因此，本发明控制Nb和V的含量之和为0～0.2％，优选地0.01～0.2％。
P：0.001％～0.100％
在钢中，P是不可避免的元素。一方面P作为固溶强化元素可以相对廉价地提高钢板的强度。另一方面，当P含量超过0.100％时，P会在晶界处偏聚，从而导致显著的韧性下降等不良影响。因此，P含量上限不大于0.100％，优选地不大于0.050％。此外，考虑到使P含量小于0.001％会增加冶炼成本，因此，P含量下限不小于0.001％，优选地不小于0.004％。
S：0.0001％～0.100％
与P相同，S同样是钢中不可避免的元素，会与钢中的Mn反应以MnS形式成为钢中的夹杂物。在S含量超过0.100％时，大量的MnS将显著损害钢的延性和韧性，使得加工性劣化。因此，S含量上限不大于0.100％，优选地不大于0.015％。同样，考虑到使S含量低于0.0001％会增加冶炼成本，因此，S含量下限不小于0.0001％，优选地不小于0.0005％，更优选地不小于0.001％。
另外，热冲压成形构件的钢板基体中可含的其他元素进行说明。
W、Mo、Ni、Cu、Co：0.01～0.30％
W、Mo、Ni、Cu、Co的添加可提升钢材的淬透性，但这几种合金元素的添加均会带来合金成本的增加，因此，在材料中仅适量添加。同时，Ni、Cu和Co还具有改善材料韧性的好处。当各元素的含量不小于0.01％时，即可表现出上述有益的效果。因此，优选地，该5种元素的含量下限分别不小于0.01％。同时，考虑材料合金的经济性，优选地，该5种元素的含量上限分别不大于0.30％。在这种情况下，在能够改善韧性的同时，确保对钢材的淬透性影响较小，保证钢材的可加工性。
Sn：0.005～0.300％，Sb：0.005～0.100％
当Sn和Sb含量不小于0.005％时能够表现出改善镀层的润湿性的效果。因此，Sn、Sb的含量下限优选地分别为不小于0.005％。不过，在含有超过0.300％的Sn和/或含有超过0.100％的Sb的情况下，会导致材料韧性变差。因此，Sn的含量优选地不大于0.300％，Sb的含量优选地不大于0.100％。
Ca：0.0001～0.01％，Mg：0.0001～0.01％，Zr：0.0001～0.01％，REM：0.0001～0.01％
不小于0.0001％的REM(Rare Earth Metal：稀土元素)能够在钢中起到净化钢液、改变钢中杂质的形态和分布、细化晶粒的作用。Ca、Mg、Zr通过分别不小于0.0001％的含量，能够实现细化夹杂物的效果，从而改善材料的性能。因此，Ca、Mg、Zr、REM的含量优选分别为不小于0.0001％。另一方面，各元素的含量超过0.01％时，上述效果饱和。因此，Ca、Mg、Zr、REM的含量分别优选为不大于0.01％。
优选地，出于考虑合金的成本，W、Mo、Ni、Cu、Co、Sn、Sb、Ca、Mg、Zr和REM元素的含量之和：在0.0001％～0.30％范围内。总体而言，本发明的钢板基体可能含有W、Mo、Ni、Cu、Co、Sn、Sb、Ca、Mg、Zr和REM等元素，这些元素的存在不影响本发明的技术问题的解决。
在本实施方式中，对于所述热冲压成形构件的钢板基体的其他成分没有特别限定。 例如，有时从废料中混入As等元素，但是如果是通常范围则不会影响所述热冲压成形构件的特性。
优选地，本发明的热冲压成形构件的第一示例钢板基体以质量百分数计包含：0.059％≤C≤0.089％、0.10％≤Si≤0.30％、1.88％≤Mn≤2.47％、0.08％≤Cr≤0.31％、0.015％≤Al≤0.05％、0.0015％≤B≤0.0035％、0.018％≤Ti≤0.046％、0≤Nb+V≤0.15％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足2.20％≤Mn+0.26Si+1.3Cr≤2.82％。优选地，第一示例钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％、0.0001％≤REM≤0.01％中的至少一者。进一步优选地，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
进一步优选地，热冲压成形构件的第一示例钢板基体以质量百分数计余量为Fe及不可避免的杂质。
更优选地，本发明的热冲压成形构件的第二示例钢板基体以质量百分数计包含：0.059％≤C≤0.089％、0.10％≤Si≤0.26％、1.88％≤Mn≤2.40％、0.11％≤Cr≤0.31％、0.015％≤Al≤0.05％、0.0015％≤B≤0.0035％、0.018％≤Ti≤0.046％、0≤Nb+V≤0.15％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足2.27％≤Mn+0.26Si+1.3Cr≤2.60％。优选地，第二示例钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％、0.0001％≤REM≤0.01％中的至少一者。进一步优选地，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
进一步优选地，热冲压成形构件的第二示例钢板基体以质量百分数计余量为Fe及不可避免的杂质。
本发明通过合金设计使得构件能够具有充足的淬透性、适当的马氏体相变温度和奥氏体转变结束温度。
在本发明中，因为低C设计使得马氏体本身具有较好的韧性，而尽可能多的马氏体组织能够实现高强度，特别是全马氏体组织可以保证构件具有更高的屈服强度，有利于阻碍构件在发生碰撞时过早发生变形失效。因此，本发明希望热冲压成形构件的室温组织尽可能为单一的马氏体组织。为此，本发明通过合理配比几种经济性的合金元素Mn、Si和Cr使 得其本身的含量有助于减少贝氏体、铁素体等非马氏体产生并且通过满足Mn+0.26Si+1.3Cr≥2.20％来使构件具备充足的淬透性。优选地，Mn+0.26Si+1.3Cr≥2.25％，更优选地Mn+0.26Si+1.3Cr≥2.50％，进一步提高淬透性，从而在更大程度上获得全马氏体组织。
本发明的构件的显微组织以面积百分数计由如下构成：小于5％的贝氏体、小于3％的奥氏体、小于3％的铁素体，剩余部分为马氏体。优选地，贝氏体、铁素体及残余奥氏体的含量总和少于5％。更优选地，贝氏体、铁素体及残余奥氏体的含量总和不超过3％。
进一步，为了保证汽车用构件的安全性，期望获得高强度的同时具备高的韧性。众所周知，在钢的热处理过程中，随着马氏体相变发生，相变潜热产生，生成的新鲜马氏体组织随着相变潜热的释放发生自回火，在消除相变应力的同时降低了马氏体中C与位错的交互作用，最终改善了马氏体(即基体)的韧性。在构件进行热变形和模具冷却的过程中，Ms温度对于控制马氏体生成及其自回火效果异常重要。本领域技术人员知晓钢中的C元素将显著降低钢的Ms温度，因此，在具有中高C含量的马氏体钢中，通常通过合金设计提高材料的Ms温度，进而提高马氏体的自回火效果，从而改善高C的硬脆马氏体组织的韧性。然而，由于本发明的钢的C含量较低，所以马氏体中碳与位错的交互作用远弱于高C含量的情况，故而马氏体自身的韧性较好，自回火对韧性的提升作用有限。同时，由于低的C含量已经导致钢材的Ms温度较高，所以马氏体的自回火程度高，容易出现马氏体回火分解软化和渗碳体的析出，导致钢材的强度和韧性显著降低。因此，在本发明的低C含量的情况下，通过采用高含量的Mn并优化Si、Cr等合金元素的配比使钢材的Ms温度控制在350～410℃，优选地，360℃～400℃的范围内，从而控制马氏体基体的自回火程度，实现材料的强度和韧性的进一步优化。
此外，C含量的降低也将显著提高Ac3温度，导致在常用的热冲压加热条件下无法获得完全奥氏体组织。为获得完全的奥氏体组织，本发明通过采用高含量的Mn并优化Si、Cr等合金元素的配比来保证Ac3温度不高于900℃。
优选地，本发明希望控制显微组织中的Nb、V、Ti等微合金碳化物析出，适量的微合金碳化物的析出将带来细化晶粒和析出强化效果，从而改善构件的强韧性。但过多的碳化物析出消耗了马氏体中的碳，从而降低了马氏体中碳和位错的交互作用，不利于构件获得高强度。因此，需要将Nb、V、Ti的含量限定在前述范围内。
本发明的热冲压成形构件，表面也可以涂覆有金属涂层，镀层可以是基于铝的合金镀层也可以基于锌的合金镀层。
需要指出的是，以上设计同样适用于制造上述热冲压成形构件的钢板。
材料的高强度有助于提升热冲压成形构件的轻量化效果。在热冲压成形构件发生碰撞失效时，构件首先是在碰撞部位发生局部塑性变形，随后是变形区域的折弯开裂。因此，材料的屈服强度成为重要的评价依据，其代表了钢材由弹性变形转变为塑性变形的临界强度，高的屈服强度可以显著推迟构件的塑性变形阶段，防止构件过早发生塑性变形，有利于零件碰撞后的维修。材料的折弯性能则定义了钢板的弯曲韧性，良好的折弯性能推迟构件在折弯变形时发生断裂失效的时间并避免早期的脆性断裂情况，因此，构件可以通过持续变形吸收能量，碰撞能量吸收表现将大幅改善。此外，构件发生折弯开裂时对应的峰值载荷越大，也意味着构件具有更佳的碰撞能量吸收表现。因此，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弯曲实验的最大弯曲角(受材料厚度影响)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，本发明人提出还需要考虑峰值力及折弯能量吸收。
常用检测材料韧性的方法是静态三点弯曲试验(即VDA弯曲试验，VDA 238-100标准)。通过试验可以得到最大弯曲角、最大弯曲角对应的峰值力、折弯能量吸收以及断裂应变等，从而反映出材料抵抗弯曲变形失效的能力。此外，室温拉伸试验(GB/T 228.1标准)是一种常见的测量材料强度的方法，屈服强度和抗拉强度反映了材料的抵抗拉伸变形失效的能力。另外，构件发生VDA弯曲时，变形区域的弯矩与钢板厚度的平方成正比关系，而折弯能量吸收W则代表了构件在整个VDA弯曲变形过程中产生弯矩所需能力的总和，与钢板厚度的平方也成正比。因此，针对不同厚度的试样，为排除板厚带来的能量差异，使用W/t2来表征单位折弯能量吸收性能。
基于上述，本发明的构件具有880～1100MPa的屈服强度YS和不低于1000MPa的抗拉强度TS，延伸率不低于5％，断裂应变ε不小于0.6。优选地，本发明的构件具有930～1050MPa的屈服强度和不低于1060MPa的抗拉强度，延伸率不低于7％，断裂应变不小于0.7。此外，本发明的构件的静态三点弯的折弯能量吸收W不小于50000mm·N，并且W与板厚t之间存在如下关系：W/t2≥2.88×104N/mm。
本发明提供了一种用于生产上述热冲压成形构件的热冲压方法，其包括：
A)提供热冲压成形用钢板的坯件或由所述热冲压成形用钢板预成形得到的预制件；
B)完全奥氏体化步骤：将坯件或预制件加热至900℃以上保温；
优选地，步骤B)在加热炉内完成，炉内最高温度在900～980℃的范围内，以达到完全奥氏体化。同时，为保证不同厚度坯料加热均匀，坯件或预制件在炉内的总加热时间T与坯件或预制件的厚度t(mm)满足：T＝t～(t×3+10)分钟；
C)热冲压成形步骤：将经加热的坯件或预制件转移到压机中进行热变形并在模具内冷却至300℃以下以获得所述热冲压成形构件，其中，热冲压成形构件从700℃冷却至Ms温度的平均冷速在未发生热变形的区域中不低于40℃/s，并且在发生热变形的区域中以更高的平均冷速进行冷却。优选地，热冲压成形构件的Ms温度满足：350℃≤Ms≤410℃。
高的平均冷速能够尽可能减少非马氏体相(如贝氏体、铁素体以及残余奥氏体等)的生成，从而保证获得的热冲压成形构件的强度和韧性。
更优选地，由于热变形会促进铁素体的形成，随热变形的增大，需要进一步提高冷却速率以充分抑制铁素体相变。所以在步骤C)中，在构件的至少一个热变形区域中的变形量不超过10％且该热变形区域的平均冷速不低于50℃/s。更优选地，在步骤C中，坯件或预制件的至少一个区域的热变形量不超过20％，且该至少一个区域在从700℃冷却至Ms温度的平均冷速不低于60℃/s。
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的说明本发明，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。
准备具有表1所示组分的钢板，相应的制造工艺如下：
钢材经如下工艺制备成为实验钢板，即：
a)炼钢：按照表1中所列成分由真空感应炉、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2h后在800℃～1200℃进行热轧，并在700℃以下进行卷取，形成热轧钢卷，并对热轧卷进行酸洗以清除热轧过程中产生的氧化皮；
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1.4～1.9mm的冷轧钢卷；
d)热镀：对T1、T2、T4、T5、T7、T8、T9、T10、T11和CT1～CT3的冷轧板进行热镀铝合金涂层以获得最终试验钢板，单面涂层重量为20～75g/m2；及对T6的冷轧板进行热镀锌合金涂层以获得最终试验钢板，单面涂层重量为40～80g/m2；
e)退火：对T3成分的冷轧板进行退火处理以获得最终试验钢板，T3未经历热镀。
表1基体钢的化学成分(wt.％，余量为Fe和不可避免的杂质)
T1～T11为在本发明的成分范围内且满足Mn、Si及Cr配比关系的示例钢，CT1～CT3 为在本发明之外的对比钢。
T1～T11的C含量在0.059～0.089％的范围内，Mn+0.26Si+1.3Cr的值不小于2.20％，Si+Cr的值不大于0.47％。T1～T11的计算得到的Ac3温度均不高于900℃，计算得到的Ms温度在366～396℃范围内。
作为对比，CT1的Si含量偏高，Mn含量偏低，使得计算得到的Mn+0.26Si+1.3Cr低于2.20，淬透性差且高Si促进铁素体生成，易生成非马氏体相，从而不利于提高强度。此外，计算得到的Ms温度也较高，达到417℃，从而无法有效控制马氏体自回火的程度，也不利于提高强度。CT2的C和Mn含量偏高，导致Ms温度较低，不能有效实现马氏体自回火以改善马氏体韧性。CT3的成分范围均在本发明的成分范围内，但是计算得到的Mn+0.26Si+1.3Cr低于2.20，表明淬透性差。这种情况下，虽然Si含量在本发明的范围内，但是淬透性不足以抑制Si对铁素体的促进，使得仍然会生成少量铁素体，不利于提高强度。此外，计算得到的Ac3温度也较高，达到906℃，不利于热冲压加热过程获得完全奥氏体。最后，CT3中同时添加了较多的Si和Cr，Si+Cr的值达到了0.52％，在制备CT3的涂镀样板时发现其存在较其他示例更为显著的漏镀以及镀层附着力不佳等表面质量问题。因此，为了进一步改善镀层钢板的涂镀表面质量，同时使得热冲压后获得更均匀的组织和更优异的钢板性能，Si+Cr的值应不超过0.50％。
临界冷速测试
采用DIL805相变仪测试T1、T2、T4、T5和CT1成分发生完全的马氏体相变临界冷却速度(以下简称“临界冷速”)。通常情下，临界冷速越大，说明材料发生完全马氏体相变所需要的冷却要求越高，即材料的淬透性越差。首先，将几种试样加热至930℃保温，在不施加变形和施加10％热变形的情况下，以20～70℃/s的冷速冷却至200℃以下，获得冷却过程(从700℃至400℃)中的膨胀曲线。判断该段膨胀曲线上是否存在非马氏体相变并记录拐点温度，以曲线上无非马氏体相变的最小冷速作为该材料的临界冷速，记录结果如表2所示。在常规热冲压成形过程中，产品的热变形量一般不超过10％，仅少量具备深拉延特征的零件，变形量可达到20％。故而，在本申请中以10％变形量作为基准来确定马氏体相变的临界冷却速度。另外，如表1中所示，根据T1和T2计算的Mn+0.26Si+1.3Cr均为2.20，相比于其他示例钢，比如T4(其Mn+0.26Si+1.3Cr为2.27)和T5(其Mn+0.26Si+1.3Cr为2.53)，T1和T2的淬透性较差。因此，在T1和T2能够满足冷却要求的情况下，其他成分也将满足要求，故而以T1和T2测试结果作为临界冷速的基准。
表2实验钢的临界冷速测试结果
由表2可知，在无热变形的情况下，T5的临界冷速在20℃/s以下，T4的临界冷速在25℃/s左右，并且T1和T2的临界冷速在30℃/s左右。值得注意的是，在热冲压成形生产中，常用的22MnB5材料的临界冷速约为27℃/s。这说明通过Mn、Si和Cr的优化设计使得Mn+0.26Si+1.3Cr的值不小于2.20％，即使C含量低，本发明的钢仍具有较好的淬透性，因此在不改变现有热冲压成形的生产条件的情况下，能够确保在冷却过程中获得全马氏体组织。
相对的，由于Mn+0.26Si+1.3Cr的值约为1.70％，在无热变形的情况下，即使在冷速达到40℃/s时，CT1在冷却过程中仍存在显著的非马氏体相变，淬透性较差。因此，在现有的热冲压成形条件下，CT1可能生成较多的铁素体和贝氏体，对生产控制以及最终零件的强度保障均不利。
在存在10％的热变形情况下，T5和T4的临界冷速分别在25～30℃/s和40～50℃/s的范围内，而T2的临界冷速在50～60℃/s的范围内。也就是说，Mn+0.26Si+1.3Cr的值越低，需要的临界冷速就越高。另外，T1和T2的Mn+0.26Si+1.3Cr的值相同，但是T1具有比T2低的临界冷速，在40～50℃/s℃的范围内。这是因为T1中添加了更多有利于抑制贝氏体生成的Cr且更少量的促进铁素体生成的Si，这有助于降低临界冷速。因此，在无热变形时淬透性接近的情况下，较高Cr含量且较低Si含量的组合添加有利于在热变形条件下使构件获得完全的马氏体组织。
相对的，CT1钢板即使冷速达到70℃/s时仍存在显著的非马氏体相变，这主要是由于CT1淬透性不足且Si含量偏高，Cr含量偏低。
鉴于上述，为获得均匀的马氏体组织，与无热变形的情况相比，存在热变形的情况 将需要更高的临界冷速。以10％变形量作为基准，对于本发明的钢材，在保证平均冷速达到60℃/s以上时，构件即可获得均匀的马氏体组织。优选地，通过添加更高含量的Cr(即0.11％以上)且较低含量的Si(即0.30％以下)能够进一步降低平均冷速至50℃/s以上，使得构件获得均匀的马氏体组织。相对的，对于CT1试样，即使平均冷却速度达到70℃/s以上，仍将生成较多的铁素体和/或贝氏体，从而导致构件无法达到最佳的性能。因此，本发明的钢材在保证性能的基础上，通过合理配比Mn、Si和Cr改善了钢材的淬透性，抑制了非马氏体相的生成，降低了对于生产控制的要求。
热冲压成形构件的性能测试
将不同厚度的T1～T11及CT1-CT3钢板试样进行热冲压帽型件生成：将尺寸为230*550mm的试样加热至910℃保温以发生完全奥氏体化，随后将试样转移至压机中进行热冲压，该转移时间为10～12s。之后，在压机中进行热冲压并保压冷却8～10s至250℃以下后取出帽型构件，冷却过程中样件的平均冷速能达到50℃/s以上。随后，在构件侧壁存在热变形的位置取样，在金相显微镜和扫描电镜下观察经历热冲压后的构件试样的显微组织，并按照GBT228.1室温拉伸标准和VDA-238三点弯曲标准进行抗拉强度、延伸率、最大弯曲角和峰值力的测试。之后，根据VDA-238三点弯曲测试中的三点弯位移-载荷曲线，计算构件试样在弯曲失效前的折弯能量吸收，即载荷达到峰值前对应的曲线的积分面积。为减小测量误差，上述测试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构件试样的VDA弯曲角αpeak；(2)基于实验结果，选取至少三组中断弯曲角αL(即构件试样在承载状态下的弯曲角)进行中断弯曲实验，保证αL≥50％αpeak；(3)当构件试样弯曲至αL时停止加载，测量构件试样在卸载状态下的弯曲角αUL；(4)将卸载构件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Ri和Ro；(5)根据等式(1)计算出不同αUL的卸载状态下的构件试样最严重变形区外表面等效(塑性)应变ε，即，构件试样弯曲至αL时构件试样最严重变形区的外表面的等效应变，从而建立ε-αL关系；(6)根据拟合结果，利用外推方法得出构件试样的弯曲断裂应变ε(即αL等于αpeak时的ε)：
最终显微组织、拉伸性能和VDA弯曲性能的最终结果见表3。
表3各构件试样的组织和性能测试结果
图1为本发明的T1构件试样的典型显微组织，其显微组织中以马氏体相为主，非马氏体的相的总量低于3％(面积百分数)。T1构件试样的典型显微组织适用于T3-T11构件试 样。图2为本发明的T2构件试样的典型显微组织，其显微组织中存在少量的贝氏体，以面积百分数计不足5％。图3为CT1构件试样的典型显微组织，其显微组织中存在较多的贝氏体，以面积百分数计达到10％以上。结合各试样构件的组织及其Mn+0.26Si+1.3Cr值，可确认当Mn+0.26Si+1.3Cr≥2.20％，能够获得以面积百分数计不足8％的非马氏体的相组成，表明了改善的淬透性，从而有利于提高钢材的强度。相对的，CT1和CT3的组织表明，当Mn+0.26Si+1.3Cr<2.20％，存在以面积百分数计超过8％的非马氏体的相组成，表明钢材的淬透性不足，不能够有效提高钢材的强度。
由表3可知，本发明的T1～T11构件试样的屈服强度在880～1041MPa的范围内，抗拉强度在1054～1186MPa的范围内，延伸率在7～8％的范围内，最大弯曲角达到88～99.7°，断裂应变在0.7以上且W/t2值不小于2.9×104N/mm，表现出优异的强度与韧性的匹配。
与T1构件试样相比，虽然CT1构件试样的最大弯曲角和断裂应变有所提升，但是屈服强度、抗拉强度、峰值力以及W/t2值均出现明显的降低，分别为816MPa、961MPa、6568N和2.84×104N/mm，不满足使用要求。这是因为CTI构件试样的组织包含10％以上的贝氏体以及少量铁素体。如前所述，贝氏体能够改进材料的韧性，因此，CT1构件试样的最大弯曲角达到104°且断裂应变为0.90。但是较多贝氏体以及少量铁素体的存在均对材料的强度不利。此外，高的Ms温度使得马氏体过度自回火，从而也导致强度降低。CT1构件试样具有较多的贝氏体和少量的铁素体组织是因为：一方面，CT1构件试样的成分不在本发明的范围内，尤其是低的Mn且Mn、Si和Cr的配比不合适(Mn+0.26Si+1.3Cr＝1.70)，导致淬透性较差，趋向于生成铁素体和贝氏体等非马氏体组织。另一方面，在淬透性较差的情况下，低的Cr含量(0.04％)不足以抑制贝氏体的生成且高的Si含量(0.42％)促进铁素体生成，从而进一步导致在相同的热冲压成形条件下，显微组织中生成较多贝氏体和铁素体。
CT2构件试样的强度较高，如抗拉强度为1267MPa，且具有与T1构件试样类似的组织。这一方面是因为Mn+0.26Si+1.3Cr＝2.94，表明CT2构件试样具有高的淬透性，能够保证全马氏体组织；另一方面是由于较高的C含量。然而，CT2构件试样的韧性比所有其他构件试样都低，最大弯曲角仅为79.3°且断裂应变为0.59。原因在于Mn含量高于2.70％，会导致出现偏析，从而降低钢材的延性和韧性。另外CT2构件试样的C含量略高且Ms较低，导致冷却过程中马氏体自回火不足，从而不能有效降低C与位错的交互作用，最终未能有效改善马氏体的韧性。这二者从而最终导致CT2构件试样韧性不足。由此可知，即使能够满足淬透性要求，但是若成分超出了本发明的范围，也不能实现兼具高强度和高韧性的 构件。
与T1构件试样相比，CT3构件试样的韧性良好，但是屈服强度降低了约7％，未能达到880MPa。这是因为虽然二者的成分均在本发明的范围内，但是CT3构件试样的Mn、Si及Cr配比未能满足Mn+0.26Si+1.3Cr≥2.20％，导致其淬透性较差，生成了相对较多的贝氏体(面积百分数约5-10％)，故而其屈服强度较低。
鉴于上述，为了获得预期的高强度和高韧性，如不小于880MPa的拉伸强度、不小于80°的最大弯曲角、不小于0.6的断裂应变且W/t2值不小于2.88×104N/mm，各合金元素不仅需要在本发明的范围内，还需要使Mn、Si及Cr满足Mn+0.26Si+1.3Cr≥2.20％，二者缺一不可。
此外，如表3所示，整体来看，在满足韧性要求的情况下，即最大弯曲角不小于80°，断裂应变ε不小于0.6、峰值力不小于7500N且折弯能量吸收W/t2≥2.88×104N/mm，随着Mn+0.26Si+1.3Cr的值增加，即淬透性进一步增加，T1～T11构件试样的强度整体呈上升趋势。尤其，T4构件试样的Mn+0.26Si+1.3Cr为2.27％，屈服强度达到932MPa。因此，当Mn+0.26Si+1.3Cr为2.25％或更大时，能够获得不小于930MPa的屈服强度。进一步，T5构件试样的Mn+0.26Si+1.3Cr为2.53％，屈服强度达到954MPa。因此，当Mn+0.26Si+1.3Cr为2.50％或更大时，能够获得不小于950MPa的屈服强度。
T1构件试样与T2构件试样的韧性接近，但是T2构件试样的屈服强度和抗拉强度明显低于T1构件试样的屈服强度和抗拉强度。这是因为虽然二者具有相近的C含量且Mn+0.26Si+1.3Cr的值相同，但是相比于T1构件试样，T2构件试样的Cr含量较低，故而不足以在冷却过程中抑制贝氏体的产生，使得所获得的组织包含少量的贝氏体(见图2)。贝氏体的存在降低了T2构件试样的屈服强度和抗拉强度。CT1同样由于Cr含量偏低导致大量贝氏体存在，从而屈服强度和抗拉强度明显降低。因此，在保证淬透性和韧性的情况下，通过提高Cr含量能够进一步提高强度。优选地0.11％≤Cr≤0.5％，且更优选地，0.12％≤Cr≤0.4％。
此外，T3和T6构件试样分别为无镀层和镀锌试样，二者的各项性能均满足使用要求。由此可知，不论是镀锌还是镀铝抑或无镀层，所获得的热冲压成形构件均能够满足预期的性能要求。另外，T1～T11的实验结果也表明了本发明的具有不同厚度的试样构件均能够满足预期的性能要求。
上述结果表明，在本发明的基体钢板成分范围内且满足Mn+0.26Si+1.3Cr≥2.20％，不论是镀锌层还是镀铝层抑或无镀层，所获得的热冲压成形构件均能够满足预期的性能要 求，兼具高强度、高韧性及高吸能能力，提高了使用安全性。另外，可以通过进一步控制合金元素的范围来进一步提高性能。
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 (22)

  1. 一种热冲压成形用钢板，所述热冲压成形用钢板的钢板基体以质量百分数计包含：0.053％≤C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％。
  2. 根据权利要求1所述的热冲压成形用钢板，所述钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％和0.0001％≤REM≤0.01％中的至少一者。
  3. 根据权利要求2所述的热冲压成形用钢板，其中，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
  4. 根据权利要求1所述的热冲压成形用钢板，所述钢板基体以质量百分数计余量为Fe及不可避免的杂质。
  5. 根据权利要求3所述的热冲压成形用钢板，所述钢板基体以质量百分数计余量为Fe及不可避免的杂质。
  6. 根据权利要求1所述的热冲压成形用钢板，其中，Mn+0.26Si+1.3Cr≥2.25％。
  7. 根据权利要求1所述的热冲压成形用钢板，其中，Mn+0.26Si+1.3Cr≥2.50％。
  8. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，0.11％≤Cr≤0.5％。
  9. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，0.12％≤Cr≤0.4％。
  10. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，2.01％≤Mn≤2.50％，和/或0.055％≤C≤0.09％且0.10％≤Si≤0.26％。
  11. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，0.01≤Nb+V≤0.2％。
  12. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，所述钢板的马氏体相变开始温度Ms通过以质量百分数计的合金元素计算为Ms＝520-320C-50Mn-5Si-20Ni-30Cr-20Mo-5Cu，且满足：350℃≤Ms≤410℃。
  13. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，所述钢板的奥氏体转变结束温度Ac3通过以质量百分数计的合金元素计算为Ac3＝912-250C-16Mn+48Si-2Cr-16Ni+95V+96Ti+210Al-10Cu，且满足：Ac3≤900℃。
  14. 根据权利要求1至7中任一项所述的热冲压成形用钢板，其中，Si+Cr≤0.47％。
  15. 一种热冲压成形构件，所述热冲压成形构件的钢板基体以质量百分数计包含：0.053％≤ C≤0.10％、0.05％≤Si≤0.30％、1.81％≤Mn≤2.7％、0.01％≤Cr≤0.7％、0.01％≤Al≤0.5％、0.0005％≤B≤0.005％、0.015％≤Ti≤0.05％、0≤Nb+V≤0.2％、0.001％≤P≤0.100％、0.0001％≤S≤0.100％，Fe≥95％及不可避免的杂质，其中，Mn、Si和Cr含量满足Mn+0.26Si+1.3Cr≥2.20％，其中，以面积百分数计，所述钢板基体的微观组织包括：小于5％的贝氏体、小于3％的奥氏体、小于3％的铁素体以及小于0.2％的Nb-V-Ti微合金碳化物，剩余部分为马氏体。
  16. 根据权利要求15所述的热冲压成形构件，所述钢板基体以质量百分数计还包含：0.01％≤W≤0.30％、0.01％≤Mo≤0.30％、0.01％≤Ni≤0.30％、0.01％≤Cu≤0.30％、0.01％≤Co≤0.30％、0.005％≤Sn≤0.30％、0.005％≤Sb≤0.100％、0.0001％≤Ca≤0.01％、0.0001％≤Mg≤0.01％、0.0001％≤Zr≤0.01％和0.0001％≤REM≤0.01％中的至少一者，其中，0.0001％≤W+Mo+Ni+Cu+Co+Sn+Sb+Ca+Mg+Zr+REM≤0.30％。
  17. 根据权利要求15所述的热冲压成形构件，所述钢板基体以质量百分数计余量为Fe及不可避免的杂质。
  18. 根据权利要求15-17任一项所述的热冲压成形构件，其中，在所述钢板基体的微观组织中，贝氏体、铁素体及残余奥氏体的总和不超过3％。
  19. 根据权利要求15-17任一项所述的热冲压成形构件，其中，所述构件具有不低于930～1050MPa的屈服强度YS和不低于1060MPa的抗拉强度TS，延伸率不低于7％，断裂应变ε不小于0.7且所述构件的静态三点弯的折弯能量吸收W与板厚t满足：W/t2≥2.88×104N/mm。
  20. 根据权利要求15-17任一项所述的热冲压成形构件，其由根据权利要求2至14中任一项所述的热冲压成形用钢板制成。
  21. 一种用于生产根据权利要求15-20中任一项所述的热冲压成形构件的热冲压方法，其包括：
    A)提供根据权利要求1至14中任一项所述的热冲压成形用钢板的坯件或由所述热冲压成形用钢板预成形得到的预制件；
    B)完全奥氏体化处理：将步骤A中的坯件或预制件加热至900℃以上并保温；
    C)热冲压成形处理：在完成步骤B之后，将经加热的坯件或预制件转移到压机中进行热变形并在模具内冷却至300℃以下以获得热冲压成形构件，其中，热冲压成形构件从700℃冷却至Ms温度的平均冷速在未发生热变形的区域中不低于40℃/s，并且在发生热变形的区域中以更高的平均冷速进行冷却。
  22. 根据权利要求21所述的方法，其中，在步骤C中，所述坯件或预制件的至少一个区域的热变形量不超过10％，且该至少一个区域从700℃冷却至Ms温度的平均冷速不低于50℃/s。
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Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN110172636A (zh) * 2018-06-28 2019-08-27 镕凝精工新材料科技(上海)有限公司 一种低碳热成形钢及其制备方法
JP2021155794A (ja) * 2020-03-26 2021-10-07 日本製鉄株式会社 ホットスタンプ部品用鋼板およびその製造方法
JP2021181618A (ja) * 2020-05-15 2021-11-25 Ｊｆｅスチール株式会社 熱間プレス部材およびその製造方法
CN114959441A (zh) * 2021-02-26 2022-08-30 宝山钢铁股份有限公司 一种热冲压钢板、热冲压部件及其制造方法
CN115261742A (zh) * 2021-04-30 2022-11-01 宝山钢铁股份有限公司 一种抗拉强度1000MPa热冲压部件及其制造方法
CN115652218A (zh) * 2022-11-17 2023-01-31 育材堂（苏州）材料科技有限公司 一种低碳的高韧性热冲压成形构件及钢板

Family Cites Families (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN104513927B (zh) * 2014-12-19 2017-04-05 宝山钢铁股份有限公司 一种抗拉强度800MPa级高强度高韧性钢板及其制造方法
WO2016198906A1 (fr) * 2015-06-10 2016-12-15 Arcelormittal Acier a haute résistance et procédé de fabrication
WO2017006144A1 (en) * 2015-07-09 2017-01-12 Arcelormittal Steel for press hardening and press hardened part manufactured from such steel
WO2018220430A1 (en) * 2017-06-02 2018-12-06 Arcelormittal Steel sheet for manufacturing press hardened parts, press hardened part having a combination of high strength and crash ductility, and manufacturing methods thereof
CN114703427A (zh) * 2018-04-28 2022-07-05 育材堂（苏州）材料科技有限公司 热冲压成形用钢材、热冲压成形工艺及热冲压成形构件
CN113106338B (zh) * 2021-03-22 2022-02-11 北京科技大学 一种超高强度高塑性热冲压成形钢的制备方法
CN114107636B (zh) * 2021-10-19 2023-02-24 北京科技大学 一种2000MPa级超高强韧轮辐用热轧热成形钢及其制备方法
CN113802065B (zh) * 2021-11-18 2022-03-29 育材堂（苏州）材料科技有限公司 热冲压成形构件、热冲压成形用钢板以及热冲压工艺
CN115058650B (zh) * 2022-06-21 2023-10-17 清华大学 高塑性热成形钢及其快速加热制备方法和应用

Patent Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN110172636A (zh) * 2018-06-28 2019-08-27 镕凝精工新材料科技(上海)有限公司 一种低碳热成形钢及其制备方法
JP2021155794A (ja) * 2020-03-26 2021-10-07 日本製鉄株式会社 ホットスタンプ部品用鋼板およびその製造方法
JP2021181618A (ja) * 2020-05-15 2021-11-25 Ｊｆｅスチール株式会社 熱間プレス部材およびその製造方法
CN114959441A (zh) * 2021-02-26 2022-08-30 宝山钢铁股份有限公司 一种热冲压钢板、热冲压部件及其制造方法
CN115261742A (zh) * 2021-04-30 2022-11-01 宝山钢铁股份有限公司 一种抗拉强度1000MPa热冲压部件及其制造方法
CN115652218A (zh) * 2022-11-17 2023-01-31 育材堂（苏州）材料科技有限公司 一种低碳的高韧性热冲压成形构件及钢板

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