Classifications

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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
  • C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
  • C23C2/36—Elongated material
  • C23C2/40—Plates; Strips
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  • 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
  • C21D1/18—Hardening; Quenching with or without subsequent tempering
  • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/005—Heat treatment of ferrous alloys containing Mn
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0226—Hot rolling
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0236—Cold rolling
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  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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  • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0273—Final recrystallisation annealing
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • 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
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
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  • 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
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  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
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  • 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
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
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  • 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
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  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
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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
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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
  • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
  • C23C2/0224—Two or more thermal pretreatments
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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
  • C23C2/04—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
  • C23C2/06—Zinc or cadmium or alloys based thereon
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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
  • C23C2/26—After-treatment
  • 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
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
  • C23C2/29—Cooling or quenching
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/001—Austenite
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/002—Bainite
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/004—Dispersions; Precipitations
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite

Definitions

• the present invention relates to a high-strength steel sheet that is applied to various usages including automobiles and home electric appliances and that has excellent formability suitable in particular for energy-absorbing members, and to a member and methods for producing the same.
• high-strength steel sheets that have a tensile strength (TS) of 980 MPa or more have been increasingly applied to frame parts and seat parts of automobiles.
• TS tensile strength
• these high-strength steel sheets are required to have formability superior to what has been typically achieved.
• Patent Literature 1 discloses that a steel sheet having a tensile strength of 980 MPa or more and excellent elongation and flangeability is obtained, the steel sheet containing, in mass%, C: 0.15% or more and 0.30% or less, P: 0.040% or less, S: 0.0100% or less, N: 0.0100% or less, O: 0.0060% or less, either or both of Si and Al: 0.70% or more and 2.50% or less in total, either or both of Mn and Cr: 1.50% or more and 3.50% or less in total, Mo: 0% or more and 1.00% or less, Ni: 0% or more and 1.00% or less, Cu: 0% or more and 1.00% or less, Nb: 0% or more and 0.30% or less, Ti: 0% or more and 0.30% or less, V: 0% or more and 0.30% or less, B: 0% or more and 0.0050% or less, Ca: 0% or more and 0.0400% or less
• Patent Literature 2 discloses that a high-strength hot-dip galvanized steel sheet having a tensile strength of 980 MPa or more and excellent fracture resistance upon collision is obtained, the steel sheet having a steel composition containing, in mass%, C: 0.07 to 0.20%, Si: 0.1 to 2.0%, Mn: 2.0 to 3.5%, P: 0.05% or less, S: 0.05% or less, and Sol.
• Al 0.005 to 0.1% with the balance being Fe and incidental impurities, and having a steel microstructure in which area fractions are as follows: ferrite: 60% or less, tempered martensite: 40% or more, and fresh martensite: 10% or less, in which the void number-density in a bent portion in a VDA bending test is adjusted to 1500 voids/mm 2 or less.
• Patent Literature 3 discloses that a high-strength cold rolled steel sheet having excellent strength, ductility, and flangeability is obtained, the steel sheet containing, in mass%, C: 0.10 to 0.40%, Mn: 0.5 to 4.0%, Si: 0.005 to 2.5%, Al: 0.005 to 2.5%, Cr: 0 to 1.0%, and the balance being iron and incidental impurities with P limited to 0.05% or less, S limited to 0.02%, and N limited to 0.006% or less, and having a steel microstructure containing, in area fractions, 2 to 30% of retained austenite and 20% or less of martensite, in which the average grain size of cementite is 0.01 ⁇ m or more and 1 ⁇ m or less, and the aforementioned cementite contains 30% or more and 100% or less of cementite having an aspect ratio of 1 or more and 3 or less.
• Patent Literature 1 and Patent Literature 3 Although a high-strength steel sheet having excellent ductility and stretch flangeability is provided, the axial crash properties are not considered.
• Patent Literature 2 discloses a high-strength steel sheet of 980 MPa or more having excellent fracture resistance upon collision; however, formability is not considered, and the technology is limited to a steel sheet having a hot-dip galvanized layer.
• the present invention has been made to address the aforementioned issues, and an object thereof is to provide a steel sheet and a member having a tensile strength (TS) of 980 MPa or more, high ductility and stretch flangeability, and excellent axial crash properties, and methods for producing the same.
• TS tensile strength
• the tensile strength is measured by a tensile test according to JIS Z 2241 (2011).
• high ductility means that the total elongation (T-El) as measured by a tensile test according to JIS Z 2241 (2011) is as follows: (A) in a TS range of 980 MPa or more and less than 1180 MPa, T-El is 14.0% or more; (B) in a TS range of 1180 MPa or more and less than 1320 MPa, T-El is 12.0% or more, and (C) in a TS range of 1320 MPa or more, T-El is 10.0% or more.
• excellent axial crash properties refer to a steel sheet having a VDA bend angle ⁇ of 80° or more in VDA bending.
• a VDA bending test is a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry, and is a 3-point bending test characterized by an extremely narrow roller spacing and a sharp punch.
• the VDA bending test uses a 60 mm ⁇ 60 mm square test specimen and involves supporting the test specimen by rollers having a roller diameter D of 30 mm and a roller spacing L of (thickness a 0 ⁇ 2) + 0.5 mm such that the bending ridge direction is parallel to the rolling direction, and then pushing a punch having a tip r of 0.4 mm into the test specimen from above at a stroke rate of 20 mm/min.
• the VDA bend angle ⁇ is an angle (° (more specifically, the unit is "°/mm” but hereinafter referred to as "°")) calculated by using equations (1) to (5) from the stroke S (mm) under the maximum load in the aforementioned bending test, and can be used as an indicator of the axial crash properties.
• the stroke S (mm) under the maximum load is the length (mm) by which the punch has moved from the start of testing to the time point when the maximum load is obtained.
• W mm p 2 + S ⁇ c 2 ⁇ c 2 [Math.
• the inventors of the present invention have conducted extensive studies on a steel sheet having high strength, excellent ductility and stretch flangeability, and excellent axial crash properties, and have arrived at the following conclusions.
• the present invention has been made based on the aforementioned findings and specifically provides the following.
• a steel sheet that has a tensile strength (TS) of 980 MPa or more and excellent ductility, stretch flangeability, and axial crash properties can be obtained.
• the steel sheet of the present invention has excellent axial crash properties and thus is suitable for use in energy-absorbing members.
• a steel sheet of the present invention includes a chemical composition containing, in mass%, C: 0.10% or more and 0.30% or less, Si: 0.5% or more and 2.0% or less, Mn: 1.5% or more and 3.0% or less, P: 0.10% or less, S: 0.020% or less, sol.
• Al 1.00% or less, N: 0.015% or less, and the balance being Fe and incidental impurities; and a steel microstructure containing, in area fractions, ferrite: 15% or less, tempered martensite: 30% or more and 80% or less, bainite: 10% or more and 40% or less, retained austenite: 5% or more, and fresh martensite: 10% or less, wherein an average C concentration in the retained austenite is 0.60 mass% or more, and, in a region within 100 ⁇ m from a steel sheet surface in a thickness direction, a content of Fe element present as carbides in the tempered martensite is 0.20 mass% or less on average.
• C is necessary for securing the tempered martensite amount and securing the retained austenite amount stable at room temperature, and, furthermore, stabilizes retained austenite by concentrating in the retained austenite and is thus a necessary element for improving the ductility.
• the C content thus needs to be 0.10% or more, and is preferably 0.12% or more and more preferably 0.15% or more.
• the C content thus needs to be 0.30% or less, and is preferably 0.28% or less and more preferably 0.25% or less.
• the Si suppresses formation of carbides in martensite and bainite, accelerates formation of retained austenite, and is an element useful for improving the stability of retained austenite.
• the Si content needs to be 0.5% or more, and is preferably 0.6% or more and more preferably 0.7% or more. Meanwhile, at a Si content exceeding 2.0%, an excessive increase in strength degrades ductility and stretch flangeability, and, in addition, the rolling load may increase during hot rolling, liquid metal embrittlement cracking may occur when welded with a zinc-coated material, and the chemical convertibility may be degraded.
• the Si content thus needs to be 2.0% or less, and is preferably 1.5% or less and more preferably 1.0% or less.
• Mn is effective for securing the specified area fractions of tempered martensite and/or bainite and for securing the strength. Mn is also an important element for increasing the retained austenite area fraction and improving the ductility since Mn decreases the Ms temperature of the retained austenite and stabilizes the retained austenite.
• the Mn content needs to be 1.5% or more, and is preferably 1.8% or more and more preferably 2.0% or more.
• the strength increases excessively, bainite transformation delays extensively, and thus ductility is degraded.
• the Mn segregation in the thickness direction becomes prominent, and there is a risk that the stability of the material quality would be degraded.
• the Mn content is thus 3.0% or less, and is preferably 2.8% or less and more preferably 2.5% or less.
• the P is an element that strengthen the steel, but a high P content degrades the spot weldability.
• the P content is 0.10% or less and is preferably 0.02% or less. It should be noted that P does not have to be contained but the P content is preferably 0.001% or more from the viewpoint of the production cost.
• the S content is an element that has an effect of improving the descalability in hot rolling and an effect of suppressing nitriding during annealing, but also has extensive adverse effects on bendability, flangeability, and spot weldability.
• the S content is to be at least 0.020% or less.
• the S content is preferably 0.0020% or less and more preferably less than 0.0010%. It should be noted that S does not have to be contained but the S content is preferably 0.0001% or more from the viewpoint of the production cost. The S content is more preferably 0.0005% or more.
• Al suppresses formation of carbides and is an element effective for accelerating formation of retained austenite. Moreover, Al is an element added as a deoxidizing agent in the steelmaking process.
• the lower limit of the sol. Al content is not particularly limited, but is preferably 0.005% or more and more preferably 0.01% or more for achieving stable deoxidization. Meanwhile, at a sol. Al content exceeding 1.00%, inclusions in the steel sheet increase, and the ductility is degraded. Thus, the sol. Al content is 1.00% or less and is preferably 0.50% or less.
• the sol. Al content is more preferably 0.25% or less and even more preferably 0.10% or less.
• N is an element that forms nitrides such as BN, AlN, and TiN in the steel and is an element that degrades the hot ductility of the steel and the surface quality. Moreover, a steel containing B has a negative effect of canceling the effects of B through formation of BN. At a N content exceeding 0.015%, the surface quality is extensively degraded. Thus, the N content is 0.015% or less and is preferably 0.010% or less. The N content is more preferably 0.005% or less and even more preferably 0.002% or less. It should be noted that N does not have to be contained but the N content is preferably 0.0001% or more and more preferably 0.001% or more from the viewpoint of the production cost.
• the balance other than what is described above is Fe and incidental impurities.
• the steel sheet of the present invention preferably has a chemical composition that contains the aforementioned basic components and the balance being iron (Fe) and incidental impurities.
• the chemical composition of the steel sheet of the present invention may contain the following optional elements as appropriate in addition to the aforementioned components.
• Ti, Nb, and V form fine precipitates during hot rolling or annealing and thereby increase the strength.
• the Ti content is preferably 0.010% or more
• the Nb content is preferably 0.020% or more
• the V content is preferably 0.020% or more.
• each content is preferably 0.100% or less, more preferably 0.080% or less, and even more preferably 0.050% or less.
• the B is an element that improves hardenability of the steel and has an advantage of facilitating formation of specified area fractions of tempered martensite and/or bainite.
• the B content is preferably 0.0005% or more and even more preferably 0.0010% or more.
• the B content is preferably 0.0050% or less, more preferably 0.0040% or less, and even more preferably 0.0025% or less.
• Cr and Cu not only serve as solid solution strengthening elements but also are elements that stabilize austenite and facilitate formation of a multi-phase microstructure during the cooling process in annealing.
• the Cr and Cu contents are each preferably 0.005% or more, more preferably 0.008% or more, and even more preferably 0.010% or more.
• the Cr and Cu contents are each preferably 0.020% or more.
• the Cr and Cu contents are each more preferably 0.04% or more. Meanwhile, at a Cr content and Cu content each exceeding 1.000%, formability of the steel sheet is degraded.
• each content is preferably 1.000% or less, more preferably 0.800% or less, and even more preferably 0.400% or less.
• the Cr and Cu contents are each preferably 0.200% or less and more preferably 0.100% or less.
• Sb and Sn are elements effective for suppressing decarburization in the region several tens of micrometers in the steel sheet surface layer formed by nitridation and oxidation of the steel sheet surface.
• the Sb and Sn contents are preferably each 0.002% or more, more preferably 0.004% or more, and even more preferably 0.006% or more.
• the Sb content and the Sn content are each preferably 0.200% or less, more preferably 0.100% or less, and even more preferably 0.040% or less.
• Ta forms alloy carbides and alloy carbonitrides and contributes to increasing the strength.
• Ta has an effect of extensively suppressing coarsening of precipitates and an effect of stabilizing the contribution ratio to the steel sheet strength improvements by precipitation strengthening when Ta partly dissolves in the Nb carbides and Nb carbonitrides and forms complex precipitates such as (Nb,Ta) and (C,N).
• the Ta content is preferably 0.005% or more.
• excessive addition of Ta saturates the precipitate stabilizing effect and increases the alloying cost.
• the Ta content is preferably 0.100% or less, more preferably 0.080% or less, and even more preferably 0.050% or less.
• Ca, Mg, and REM are elements used in deoxidation, spheroidize the sulfides, and are elements effective for addressing adverse effects of the sulfides on the local ductility and the stretch flangeability.
• the Ca content is preferably 0.0001% or more
• the Mg content is preferably 0.0001% or more
• the REM content is preferably 0.0001% or more.
• excessive addition of Ca, Mg, and REM exceeding 0.0050% induces the increase of inclusions etc., and causes defects and the like on the surface or in the inside.
• the Ca, Mg, and REM contents are each preferably 0.0050% or less, more preferably 0.0025% or less, and even more preferably 0.0010% or less.
• the Ca content is more preferably 0.0008% or less.
• the Mg content is more preferably 0.0008% or less.
• the REM content is more preferably 0.0008% or less.
• REM in the present invention refers to scandium (Sc) having an atomic number 21, yttrium (Y) having an atomic number 39, and lanthanoids from lanthanum (La) having an atomic number 57 to lutetium (Lu) having an atomic number 71.
• a REM content in the present invention is the total content of one or more of the elements selected from among the REMs described above. REM is not particularly limited but is preferably La and/or Ce.
• the optional elements contained in amount lower than the lower limit values do not impair the effects of the present invention.
• the optional elements are contained in amounts lower than the lower limit values, the optional elements are considered incidental impurities.
• Ferrite that is formed during annealing or the cooling process contributes to improving ductility; however, a difference in hardness occurs relative to the surrounding hard phases such as martensite, cracks propagate from the interface between ferrite and the hard phases during bending, and thus the stretch flangeability and the axial crash properties are degraded. Ferrite does not impair the effects of the present invention as long as the area fraction thereof is 15% or less, and thus the ferrite area fraction is 15% or less.
• the ferrite area fraction is preferably 10% or less and more preferably 5% or less. Furthermore, the ferrite area fraction may be 0%.
• the area fraction of tempered martensite is 30% or more.
• the area fraction of tempered martensite is preferably 40% or more and more preferably 50% or more. Meanwhile, at a tempered martensite area fraction exceeding 80%, the ductility is degraded due to the excessive increase in strength, and thus the tempered martensite area fraction is 80% or less, preferably 75% or less, and more preferably 70% or less.
• Bainite in the present invention is a general term for upper bainite and lower bainite.
• upper bainite and lower bainite are defined as follows.
• Upper bainite is composed of lath-like bainitic ferrite, and retained austenite and/or carbides present between laths of bainitic ferrite and is characterized by the absence of fine carbides regularly aligned within the lath-like bainitic ferrite.
• lower bainite is composed of lath-like bainitic ferrite and retained austenite and/or carbides present between laths of bainitic ferrite, but lower bainite is characterized by the presence of fine carbides regularly aligned within the lath-like bainitic ferrite.
• upper bainite and lower bainite are distinguished from each other by the absence or presence of fine carbides that are regularly aligned within bainitic ferrite.
• Bainite that is formed during the cooling process after annealing and the retention process in the cooling process causes carbon to concentrate to the surrounding non-transformed austenite region and renders it possible to efficiently obtain more stable retained austenite.
• the bainite area fraction is 10% or more.
• bainite is soft compared with tempered martensite, excessive formation of bainite decreases the strength and degrades the stretch flangeability due to the increase in the difference in hardness between bainite and tempered martensite.
• the bainite area fraction is 40% or less, preferably 35% or less, and more preferably 30% or less.
• upper bainite causes concentration of a larger amount of carbon to the surrounding non-transformed austenite region, and thus is advantageous from the viewpoint of efficiently obtaining retained austenite.
• it is particularly preferable to contain 10% or more and more preferably 15% or more of upper bainite.
• upper bainite is soft compared with lower bainite, excessive formation of upper bainite is disadvantageous from the viewpoint of improving strength and stretch flangeability.
• the upper bainite area fraction is preferably 30% or less and more preferably 25% or less.
• the retained austenite area fraction relative to the entire steel microstructure is 5% or more.
• the retained austenite area fraction is preferably 7% or more and more preferably 10% or more.
• This amount of retained austenite includes retained austenite present between bainitic ferrite.
• the area fraction of the retained austenite present between bainitic ferrite can be measured in the same manner as area fractions of other retained austenite according to the measurement method described below. Meanwhile, when the amount of retained austenite increases excessively, the strength may decrease, and the stretch flangeability and the delayed fracture resistance may be degraded.
• the retained austenite area fraction is preferably 20% or less and more preferably 15% or less.
• the fresh martensite area fraction is 10% or less, preferably 8% or less, and more preferably 5% or less since fresh martensite may cause degradation of ductility and stretch flangeability.
• the fresh martensite area fraction may be 0%.
• the average C concentration in the retained austenite is 0.60 mass% or more and is preferably 0.80 mass% or more.
• the upper limit is not particularly limited, but when the average C concentration in the retained austenite is excessively increased, the retained austenite does not undergo transformation during deformation, and the ductility improving effect may not be sufficiently obtained; thus the average C concentration in the retained austenite is preferably 2.0 mass% or less and more preferably 1.5 mass% or less.
• Carbides that precipitate in tempered martensite act as starting points of cracks during bending deformation and cause degradation of axial crash properties.
• precipitation of carbides is a phenomenon that competes with carbon partitioning, and since carbon is consumed by precipitation of the carbides, concentration of carbon into non-transformed austenite may be inhibited.
• the content of the Fe element present as carbides in the tempered martensite is preferably 0.20 mass% or less and more preferably 0.15 mass% or less on average. Although the lower limit is not particularly specified, in the present invention, the content of the Fe element is substantially 0.001 mass% or more.
• the steel microstructure of the present invention contains, as described above, ferrite (including 0%), tempered martensite, bainite, retained austenite, and fresh martensite (including 0%).
• the steel microstructure of the present invention may include ferrite (including 0%), tempered martensite, bainite, retained austenite, and fresh martensite (including 0%).
• the steel microstructure of the present invention may contain 5% or less of pearlite and cementite as phases other than those described above, and this does not impair the effects of the present invention.
• Fig. 1 illustrates one example of a SEM image of a steel microstructure of the steel sheet. As illustrated in Fig. 1 , ferrite (in Fig. 1 ), ferrite (in Fig.
• Fig. 1 see reference sign F
• Tempered martensite in Fig. 1 , see reference sign TM
• MA fresh martensite and/or retained austenite
• Fig. 1 see reference signs FM and RA
• bainite is defined as the sum of the area fractions of upper bainite and lower bainite.
• Upper bainite is determined by observation by SEM as with ferrite etc., described above.
• a microstructure composed of bainitic ferrite which is a region that does not involve precipitation of regularly aligned fine carbides and appears in the darkest black under SEM, and carbides or retained austenite that appears white under SEM is assumed to be upper bainite (in Fig. 1 , see reference sign UB).
• a ferrite region with aspect ratio ⁇ 2.0 is assumed to be ferrite and a region with aspect ratio > 2.0 is assumed to be bainitic ferrite.
• FIG. 2 is a diagram illustrating a method for measuring the steel microstructure of the steel sheet of the present invention.
• the aspect ratio is determined as a/b irrespective of whether the shape is an elliptic shape illustrated in (1) or a polygonal shape illustrated in (2).
• a is the length of the longest segment among the segments formed by any two points on the region perimeter of each phase
• b is the length of the segment that is perpendicular to the straight line forming a and that is the longest among the segments formed by any two points on the region perimeter.
• measurement is performed without considering the grains as one grain but separately as two or more grains.
• the carbide form of lower bainite is a single variant in which precipitation occurs regularly in one direction inside the substructure
• the carbides of tempered martensite are a multivariant in which precipitation directions are random inside the substructure.
• the area fractions of lower bainite and tempered martensite having such characteristics can be determined by observing ten 1.5 ⁇ m square fields of view by TEM, calculating the area fractions of the constituent phases (lower bainite and tempered martensite) by using Adobe Photoshop described above from ten fields of view, and averaging the obtained values.
• the area fraction of retained austenite and the average C concentration in retained austenite are determined by grinding and polishing the steel sheet in the thickness direction down to a thickness 1/4 position and performing X-ray diffraction intensity measurement thereon.
• the volume fraction of retained austenite is calculated by using a Mo tubular radiation source from the intensity ratios of the (200), (220), and (311) planes of austenite relative to the diffraction intensities of the (200) and (211) planes of ferrite.
• the average C concentration in retained austenite is determined by using a Co tubular radiation source, determining the lattice constant A of austenite from the peak angle of the (220) plane of austenite, and calculating the concentration from equation (6) below.
• Mn%, Si%, and Al% each indicate mass% of the content of each element (Mn, Si, and Al) in the steel sheet.
• Average C concentration in retained austenite (mass%) A - ⁇ (0.3572 + 0.0012 ⁇ Mn% - 0.00157 ⁇ Si% + 0.0056 ⁇ Al%) ⁇ /0.033
• the area fraction of fresh martensite is determined from equation (7) below, and is a value obtained by subtracting the area fraction of retained austenite (RA) from the area fraction of the aforementioned MA (fresh martensite and/or retained austenite).
• RA retained austenite
• the "area fraction" of RA (retained austenite) can be regarded as equivalent to the "volume fraction" of retained austenite determined by XRD measurement.
• FM % MA % ⁇ RA %
• the average content of the Fe element present as carbides in tempered martensite is determined by extracted residue analysis.
• a 10 mass% AA electrolyte acetylacetone-1 mass% tetramethylammonium chloride-methanol
• electrolysis is performed under conditions of current density of 20 mA/cm 2 and an electrolysis time of 30 min to dissolve a region within 100 ⁇ m from the steel sheet surface.
• the specimen is immersed in methanol, the residue separated by ultrasonic agitation (precipitates that remain undissolved by electrolysis) is filtered out by using a filter having a 0.2 ⁇ m pore diameter, and the Fe element amount in the collected residue is determined by ICP atomic emission spectroscopy and is converted into a mass% in the steel sheet.
• the collected residue is solely composed of precipitates that are present in the region within 100 ⁇ m from the steel sheet surface, and, here, the number of carbides outside the tempered martensite is significantly small (bainite, ferrite, etc.) compared with the number of carbides in the tempered martensite.
• the Fe element does not exist as iron-based compounds other than carbon-based compounds.
• the average content of the Fe element present as carbides in tempered martensite in the region within 100 ⁇ m from the steel sheet surface in the thickness direction is obtained.
• the steel sheet of the present invention described heretofore may be a steel sheet that has a zinc coating layer on a steel sheet surface (one surface or both surfaces).
• the coating layer may be a hot-dip coating layer or an electroplating layer.
• the thickness of the steel sheet of the present invention is preferably 0.5 mm or more.
• the thickness is preferably 2.0 mm or less.
• temperatures specified in each step of the present invention are surface temperatures of slabs (steel slabs) or steel sheets.
• Fig. 3 is a diagram illustrating the method for producing the steel sheet of the present invention, in particular, changes with time in surface temperature of a slab (steel slab) or a steel sheet. Details of the individual steps including the changes with time in temperature will now be described.
• Fig. 3(a) illustrates the changes with time in surface temperature of the slab (steel slab) or the steel sheet in the method for producing the steel sheet according to the first embodiment (the case in which the retention treatment is not performed).
• Fig. 3(b) illustrates the changes with time in surface temperature of the slab (steel slab) or the steel sheet in the method for producing the steel sheet according to the second embodiment (the case in which the retention treatment is performed).
• a steel slab having the aforementioned chemical composition is hot-rolled and cold-rolled, and the obtained cold rolled steel sheet is subjected to an annealing treatment at an annealing temperature of (A c3 temperature (°C) - 15°C) or higher and 900°C or lower, is cooled at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less in a range from the annealing temperature to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower, is heated at an average heating rate of 2°C/s or more in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower, is cooled at an average cooling rate CR2 of 0.5°C/s or more in a temperature range from the temperature T2 to a temperature T3
• Hot rolling may be performed according to a standard method; for example, a steel slab having the aforementioned specified chemical composition may be heated to a slab heating temperature of 1100°C or higher, hot-rolled by setting the soaking time to 20 min or longer and a finish rolling delivery temperature to the Ar 3 transformation temperature or higher, and coiled at a coiling temperature of 400°C or higher.
• the slab heating temperature may be 1300°C or lower.
• the soaking time may be 300 min or shorter.
• the finish rolling delivery temperature may be equal to or lower than Ar 3 transformation temperature + 200°C.
• the coiling temperature may be 720°C or lower.
• the coiling temperature is preferably controlled from the viewpoints of suppressing the thickness variation and stably securing high strength. Specifically, the coiling temperature is preferably 430°C or higher.
• the coiling temperature is preferably 530°C or lower.
• Cold rolling may be performed at a rolling reduction (cumulative rolling reduction) of 30% or more.
• the reduction may be 85% or less.
• the rolling reduction is preferably controlled from the viewpoints of stably obtaining high strength and decreasing anisotropy. Specifically, the rolling reduction is preferably 35% or more. Note that when the rolling load is high, it is possible to perform an annealing treatment for softening in a continuous annealing line (CAL) or a box annealing furnace (BAF) at 450°C or higher and 730°C or lower.
• CAL continuous annealing line
• BAF box annealing furnace
• the steel slab having the specified chemical composition is hot-rolled and cold-rolled and then annealed under the conditions specified below.
• the annealing facility is not particularly limited; however, from the viewpoints of productivity and achieving the desired heating rates and cooling rates, a continuous annealing line (CAL) or a continuous galvanizing line (CGL) is preferably used.
• CAL continuous annealing line
• CGL continuous galvanizing line
• the annealing temperature is A c3 temperature (°C) - 15°C or higher and 900°C or lower.
• the annealing temperature is preferably adjusted so that annealing occurs in the austenite single-phase range.
• the annealing temperature is preferably the A c3 temperature or higher.
• the annealing temperature is more preferably equal to or higher than A c3 temperature (°C) + 15°C.
• the annealing temperature is 900°C or lower.
• the annealing temperature is preferably 880°C or lower.
• the A c3 temperature can be determined by measuring the change in volume observed when a columnar test specimen (diameter of 3 mm ⁇ height of 10 mm) is heated from room temperature (25°C) to the austenite single-phase range by using a Formastor tester.
• First cooling treatment step cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less in a range from the annealing temperature to a temperature T1 of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower]
• the average cooling rate CR1 in the temperature range from the annealing temperature to a cooling stop temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower is 3°C/s or more.
• the average cooling rate CR1 is preferably 5°C/s or more and more preferably 8°C/s or more.
• the average cooling rate CR1 is 100°C/s or less.
• the average cooling rate CR1 is preferably 50°C/s or less.
• the temperature T1 (cooling stop temperature T1) is 200°C or higher.
• the temperature T1 (cooling stop temperature T1) is preferably 220°C or higher and more preferably 240°C or higher.
• the cooling stop temperature T1 is (Ms temperature (°C) - 30°C) or lower.
• the cooling stop temperature T1 is preferably (Ms temperature (°C) - 35°C) or lower.
• the martensite transformation start temperature Ms (°C) can be determined with a Formastor tester by measuring the change in volume of a columnar test specimen, which is similar to the one used in measuring the A c3 temperature, observed when the test specimen is held at a specified annealing temperature and then rapidly cooled with helium gas.
• the average cooling rate CR1 is "(annealing temperature (cooling start temperature) (°C) - (Ms temperature (°C) - 30°C) (cooling stop temperature (°C)))/(cooling time (s) from annealing temperature to (Ms temperature (°C) - 30°C))".
• Heat treatment step heating at an average heating rate of 2°C/s or more from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower]
• the temperature T2 is 300°C or higher and 450°C or lower.
• the temperature T2 is preferably 330°C or higher and more preferably 350°C or higher.
• the temperature T2 is preferably 420°C or lower and more preferably 400°C or lower.
• the average heating rate in the temperature range from the temperature T1 (cooling stop temperature T1) to a temperature T2 of 300°C or higher and 450°C or lower is 2°C/s or more.
• the average heating rate is preferably 5°C/s or more and more preferably 10°C/s or more.
• the upper limit of the average heating rate is not particularly limited, the average heating rate is preferably 50°C/s or less and more preferably 30°C/s or less.
• the average heating rate is "(temperature T2 (heating stop temperature) (°C) - temperature T1 (heating start temperature) (°C))/(heating time (s) from temperature T1 to temperature T2)".
• a second cooling treatment step that involves specified conditions is performed between the heat treatment step and the third cooling treatment step.
• the temperature T3 (cooling stop temperature T3) in the second cooling treatment step is (T2 - 150°C) or higher and (T2 - 30°C) or lower.
• the temperature T3 is preferably (T2 - 120°C) or higher and more preferably (T2 - 100°C) or higher.
• the temperature T3 is preferably (T2 - 50°C) or lower and more preferably (T2 - 70°C) or lower.
• the average cooling rate CR2 is 0.5°C/s or more, preferably 0.6°C/s or more and more preferably 0.7°C/s or more.
• the upper limit is not particularly specified, a higher average cooling rate delays carbon partitioning to retained austenite and tempering of fresh martensite, and this tends to degrade ductility and stretch flangeability; thus, the average cooling rate CR2 is preferably 5°C/s or less.
• the average cooling rate CR2 is "(temperature T2 (°C) (cooling start temperature) - temperature T3 (°C) (cooling stop temperature))/(cooling time (s) from temperature T2 to temperature T3)".
• the temperature T4 is 150°C or higher and 350°C or lower.
• the temperature T4 is preferably 180°C or higher and more preferably 200°C or higher.
• the temperature T4 is preferably 300°C or lower and more preferably 250°C or lower.
• the average cooling rate CR3 is 0.1°C/s or more, preferably 0.2°C/s or more, and more preferably 0.3°C/s or more. Although the upper limit is not particularly specified, from the viewpoint of sufficiently accelerating carbon partitioning to non-transformed austenite, the average cooling rate CR3 is preferably average cooling rate CR2 or less and 1°C/s or less.
• the average cooling rate CR3 is "(temperature T3 (°C) (cooling start temperature) - temperature T4 (°C) (cooling stop temperature))/(cooling time (s) from temperature T3 to temperature T4)".
• the holding temperature is the aforementioned temperature T4 or lower.
• the holding temperature is 150°C or higher and 350°C or lower and the temperature T4 or lower.
• the holding temperature is preferably 180°C or higher and more preferably 200°C or higher.
• the holding temperature is preferably 300°C or lower and more preferably 250°C or lower.
• the holding time is 20 s or longer and 1000 s or shorter.
• the holding time is preferably 100 s or longer and more preferably 200 s or longer.
• the holding time is preferably 700 s or shorter and more preferably 500 s or shorter.
• a zinc coating layer on a steel sheet surface by performing a hot-dip galvanizing treatment and even an alloying treatment after the first cooling treatment step (after cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less), that is, before, after, or during any one of the steps from the heat treatment step of heating from the temperature T1 to the temperature T2 to the holding treatment step of holding the temperature at 150°C or higher and 350°C or lower.
• the steel sheet is preferably dipped in a zinc coating bath at 440°C or higher and 500°C or lower to perform a hot-dip galvanizing treatment, and then the coating weight is preferably adjusted by gas wiping or the like.
• the galvanization preferably uses a zinc coating bath having an Al content of 0.10 mass% or more and 0.22 mass% or less. Moreover, an alloying treatment of the zinc coating can be performed after the hot-dip galvanizing treatment. The alloying treatment on the zinc coating is preferably performed in the temperature range of 470°C or higher and 590°C or lower, and heating to this temperature range does not impair the effects of the present invention and thus may be performed.
• the steel sheet may be skin-pass rolled.
• the skin-pass elongation is preferably 0.1% or more and 0.5% or less.
• the sheet shape can be flattened by using a leveler.
• the average cooling rate CR4 from the aforementioned holding end temperature to a temperature of 50°C or lower is 1°C/s or more and preferably 5°C/s or more.
• the upper limit of the average cooling rate CR4 is not particularly specified; however, from the viewpoint of production and the viewpoint of suppressing transformation from retained austenite to fresh martensite, the average cooling rate CR4 is preferably 10°C/s or less.
• the average cooling rate CR4 is "(150°C or higher and 350°C or lower and temperature T4 (°C) or lower (holding end temperature (cooling start temperature)) - temperature of 50°C or lower (cooling stop temperature))/(cooling time (s) from holding end temperature to cooling stop temperature)".
• a low-temperature heat treatment can be performed at 100 to 300°C for 30 s to 10 days.
• hydrogen that has penetrated the steel sheet during tempering of the martensite formed during the final cooling or skin-pass rolling and annealing is removed from the steel sheet.
• the low-temperature heat treatment can decrease hydrogen to less than 0.1 mass ppm.
• the steel sheet may be electrogalvanized after the aforementioned fourth cooling treatment step (after cooling at an average cooling rate CR4 of 1°C/s or more). From the viewpoint of decreasing hydrogen in the steel, the aforementioned low-temperature heat treatment is preferably performed after electroplating.
• a steel slab having the aforementioned chemical composition is hot-rolled and cold-rolled, and the obtained cold rolled steel sheet is subjected to an annealing treatment at an annealing temperature of (A c3 temperature (°C) - 15°C) or higher and 900°C or lower, is cooled at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less in the temperature range from the annealing temperature to 500°C (pre-retention treatment first cooling treatment), is retained for 10 s or longer and 60 s or shorter at an average cooling rate CR6 of 10°C/s or less in the temperature range from 500°C to a retention stop temperature T5 of the Ms temperature (°C) or higher and 320°C or higher (retention treatment), is cooled at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less in
• hot rolling, cold rolling, and the annealing treatment can be performed under the same conditions as those in the first embodiment.
• the treatment in the first cooling treatment step of the first embodiment is replaced by the pre-retention treatment first cooling treatment, the retention treatment, and the post-retention treatment first cooling treatment.
• the heat treatment, the second cooling treatment, the third cooling treatment, the holding treatment, and the fourth cooling treatment after the post-retention treatment first cooling treatment can be performed under the same conditions as those of the heat treatment, the second cooling treatment, the third cooling treatment, the holding treatment, and the fourth cooling treatment of the first embodiment, respectively.
• the hot-dip galvanizing treatment in the second embodiment can be performed under the same conditions as those of the first embodiment except that, whereas the hot-dip galvanizing treatment is performed after the first cooling treatment step (after cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less) in the first embodiment, the hot-dip galvanizing treatment is performed after the post-retention treatment first cooling treatment step (after cooling at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less) in the second embodiment.
• other conditions including those of the low-temperature heat treatment after all of the heat treatments and the electroplating treatment can be the same as those of the first embodiment.
• the pre-retention treatment first cooling treatment, the retention treatment, and the post-retention treatment first cooling treatment are mainly described below.
• Pre-retention treatment first cooling treatment step cooling at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less in a temperature range from the annealing temperature to 500°C]
• Retention treatment retaining at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter in a temperature range from 500°C to a retention stop temperature T5 of the Ms temperature (°C) or higher and 320°C or higher]
• the retention treatment By performing the retention treatment during cooling, a larger area fraction of bainite is obtained, carbon concentrates into non-transformed austenite from this bainite, the retained austenite becomes stable, and thus a steel sheet having more superior ductility can be obtained.
• the retention treatment is preferably performed during cooling.
• the temperature range from the annealing temperature to 500°C is cooled at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less.
• the average cooling rate CRS is preferably 8°C/s or more.
• the average cooling rate CR5 is 100°C/s or less, preferably 50°C/s or less, and more preferably less than 30°C/s.
• the average cooling rate CR5 is "(annealing temperature (cooling start temperature) (°C) - 500°C (cooling stop temperature))/cooling time (s) from annealing temperature to 500°C".
• the temperature range from 500°C to a retention stop temperature T5 of the martensite transformation start temperature Ms (°C) or higher and 320°C or higher is retained at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter; in this manner, bainite can be formed, and retained austenite having a high C concentration can be formed adjacent thereto.
• Ms temperature (°C) martensite temperature
• bainite transformation progresses excessively due to the swing back phenomenon, and the strength and the stretch flangeability are degraded.
• the aforementioned temperature range exceeds 500°C the drive force of the bainite transformation decreases, and the bainite transformation amount decreases.
• the aforementioned temperature range is lower than 320°C, lower bainite is mainly formed; however, from the viewpoint of improving ductility and VDA bending properties, upper bainite that does not accompany carbide precipitation is advantageous from the viewpoint of suppressing carbon-concentrated carbides in the non-transformed ⁇ caused by carbide precipitation, and of degradation of the axial crash properties.
• the retention stop temperature T5 is 320°C or higher.
• the temperature range is the Ms temperature (°C) or higher and 320°C or higher and 500°C or lower.
• the aforementioned temperature range is preferably 350°C or higher and more preferably 380°C or higher.
• the aforementioned temperature range is preferably 480°C or lower and more preferably 460°C or lower.
• the bainite transformation amount decreases when the average cooling rate CR6 exceeds 10°C/s.
• the average cooling rate CR6 is 10°C/s or less.
• the retention time is 10 s or longer and 60 s or shorter.
• the retention time is preferably 20 s or longer.
• the retention time is preferably 50 s or shorter.
• the average cooling rate CR6 is "(500°C (retention start temperature) (°C) - (retention stop temperature (°C) of Ms temperature (°C) or higher and 320°C or higher))/(retention time (s) from retention start temperature to retention stop temperature)".
• Post-retention treatment first cooling treatment step cooling at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less in a range from the retention stop temperature T5 to a temperature T1 of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower]
• the average cooling rate CR7 in the temperature range from the retention stop temperature T5 to the cooling stop temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower is 3°C/s or more.
• the average cooling rate CR7 is preferably 5°C/s or more and more preferably 8°C/s or more.
• the average cooling rate CR7 is 100°C/s or less.
• the average cooling rate CR7 is preferably 50°C/s or less.
• the temperature T1 (cooling stop temperature T1) is 200°C or higher.
• the temperature T1 (cooling stop temperature T1) is preferably 220°C or higher and more preferably 240°C or higher.
• the cooling stop temperature T1 is (Ms temperature (°C) - 30°C) or lower.
• the cooling stop temperature T1 is preferably (Ms temperature (°C) - 35°C) or lower.
• the average cooling rate CR7 is "(retention stop temperature T5 (cooling start temperature) (°C) - (Ms temperature (°C) - 30°C) (cooling stop temperature (°C)))/(cooling time (s) from retention stop temperature T5 to cooling stop temperature)".
• a member of the present invention is obtained by subjecting the steel sheet of the present invention to at least one of a forming process or a joining process. Furthermore, a method for producing a member according to the present invention includes a step of subjecting the steel sheet of the present invention to at least one of a forming process or a joining process to obtain a member.
• the steel sheet of the present invention has a tensile strength of 980 MPa or more and has high ductility, excellent stretch flangeability, and excellent axial crash properties.
• the member obtained by using the steel sheet of the present invention also has a high strength and has high ductility, excellent stretch flangeability, and excellent axial crash properties compared with typical high-strength members.
• weight reduction is possible by using the member of the present invention.
• the member of the present invention is suitable for use in automotive frame parts.
• the member of the present invention includes a welded joint.
• the forming process may involve a common processing method such as pressing without any limitation.
• the joining process may involve a common welding such as spot welding or arc welding, riveting, crimping, or the like without any limitation.
• Cold rolled steel sheets having a thickness of 1.4 mm and chemical compositions indicated in Table 1 were treated under the annealing conditions indicated in Tables 2 and 3 to produce steel sheets of the present invention and steel sheets of comparative examples.
• the cold rolled steel sheets were obtained by subjecting steel slabs having chemical compositions indicated in Table 1 to hot rolling (heating temperature: 1250°C, soaking time: 60 min, finish rolling delivery temperature: 1150°C, coiling temperature: 550°C) and cold rolling (rolling reduction (rolling cumulative reduction): 50%).
• Table 2 indicates conditions that do not involve the retention treatment
• Table 3 indicates conditions that involve the retention treatment.
• a hot-dip galvanizing treatment was performed in one of the steps after heating from the temperature T1 to holding a temperature of 150°C or higher and 350°C or lower to prepare hot-dip galvanized steel sheets (GI).
• the steel sheet was dipped in a zinc coating bath at 440°C or higher and 500°C or lower to perform a hot-dip galvanizing treatment, and then the coating weight was adjusted by gas wiping or the like.
• the hot-dip galvanization used a zinc coating bath having an Al content of 0.10% or more and 0.22% or less.
• hot-dip galvanized steel sheets were alloyed after the hot-dip galvanizing treatment into hot-dip galvannealed steel sheets (GA).
• the alloying treatment was conducted in the temperature range of 460°C or higher and 590°C or lower.
• some steel sheets (cold rolled steel sheets: CR) were electroplated into electrogalvanized steel sheets (EG). [Table 2] No. Steel No.
• Tables 4 and 5 The steel microstructure was measured by the aforementioned method. The measurement results are shown in Tables 4 and 5. Table 4 indicates conditions that do not involve the retention treatment, and Table 5 indicates conditions that involve the retention treatment.
• samples having a T-El of 14.0% or more with TS of less than 1180 MPa, samples having a T-El of 12.0% or more with TS of 1180 MPa or more and less than 1320 MPa, and samples having a T-El of 10.0% or more with TS of 1320 MPa or more were determined as having excellent ductility.
• the stretch flangeability was evaluated by a hole expansion test according to the Japan Iron and Steel Federation Standard, JFS T1001. That is, after a 100 mm ⁇ 100 mm square sample was punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance), the sample was arranged such that the burr, which had been formed by making the hole by punching, to face outward, and the hole was expanded with a conical punch having a 60 degree apex angle until cracks that penetrated through the thickness occurred.
• JFS T1001 Japan Iron and Steel Federation Standard
• the hole expansion ratios ⁇ are indicated in Tables 4 and 5. Steel sheets having ⁇ of 45% or more were determined as having excellent stretch flangeability.
• Fig. 4 is a diagram illustrating a method for calculating the VDA bend angle.
• the VDA bend angle ⁇ was evaluated as the evaluation of the axial crash properties by a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry.
• the VDA bending test is a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry, and is a 3-point bending test characterized by rollers 10 with an extremely narrow roller spacing and a sharp punch 11.
• the VDA bend angle ⁇ is an angle (°) calculated by using equations (1) to (5) from the stroke S (mm) under the maximum load in the aforementioned bending test, and can be used as an indicator of the axial crash properties.
• Steel sheets having ⁇ of 80° or more were determined as having excellent axial crash properties.
• the examples of the present invention indicated in Tables 4 and 5 are excellent in all of strength, ductility, stretch flangeability, and axial crash properties, but the comparative examples are poor in at least one.
• a member obtained by forming a steel sheet of the example of the present invention a member obtained by joining the steel sheet, and a member obtained by forming and joining the steel sheet had high strength, high ductility, excellent stretch flangeability, and excellent axial crash properties as with the steel sheet of the example of the present invention since the steel sheet of the example of the present invention had high strength, high ductility, excellent stretch flangeability, and excellent axial crash properties.
• the present invention enables production of a steel sheet that has excellent ductility, stretch flangeability, and axial crash properties, that is used in members of automobiles and home electric appliances, and that is particularly suitable for use in energy-absorbing members of automobiles, and such a steel sheet is suitable for press-forming of these members.

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