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
  • C23C2/04—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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  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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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
  • 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
  • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
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  • 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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  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
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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 by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties 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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  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties 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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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0263—Modifying the physical properties 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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  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0273—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
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  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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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
  • 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
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
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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
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
  • C23C2/29—Cooling or quenching
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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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  • 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
  • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
  • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
  • C23G1/02—Cleaning or pickling metallic material with solutions or molten salts with acid solutions
  • C23G1/08—Iron or steel
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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
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/001—Austenite
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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
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/002—Bainite
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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
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite
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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
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C18/00—Alloys based on zinc

Definitions

• the present invention relates to a steel sheet, a member, and methods for producing the same.
• Automotive steel sheets are required to have strength and excellent stretch flangeability. In addition, high toughness is necessary so that the automotive parts do not fracture upon collision.
• Patent Literature 1 discloses a high-strength steel sheet having good workability and a method for producing the same.
• Patent Literature 2 discloses a high-strength cold rolled steel sheet and a method for producing the same.
• Patent Literature 1 and Patent Literature 2 all have a tensile strength TS of less than 1180 MPa, and the stretch flangeability and the toughness are not considered.
• the present invention has been made under such circumstances, and an object thereof is to provide a high-strength steel sheet and a member having a tensile strength TS of 1180 MPa or more and excellent stretch flangeability and toughness, and methods for producing the same.
• high strength means that the tensile strength TS as measured in conformity with JIS Z 2241 (2011) is 1180 MPa or more.
• excellent stretch flangeability means that the hole expansion ratio as measured in conformity with JIS Z 2256 (2010) is 30% or more.
• excellent toughness means that, in a Charpy impact test performed in conformity with JIS Z 2242 (2016), the brittle-to-ductile transition temperature is - 40°C or lower.
• the inventors of the present invention have earnestly conducted studies to achieve the aforementioned object and thus gained following findings.
• the main phase is preferably martensite or bainite.
• the area fraction of ferrite is 1% or less, the total area fraction of martensite and bainite is 95% or more, and the area fraction of the retained austenite is 5% or less, void formation at the interfaces between soft ferrite and hard martensite or bainite is suppressed, and the stretch flangeability is improved.
• addition of Nb decreases the prior-austenite grain size to 10 ⁇ m or less, and the toughness is improved.
• addition of B causes B to segregate to the prior-austenite grain boundaries and strengthen the grain boundaries, thereby improving the toughness.
• some cases have seen moderate improvements in toughness and while others have seen significant improvements in toughness; thus, conditions that have caused the toughness to significantly improve have been investigated in further detail.
• an austenite microstructure having the same crystal orientation as in the first annealing is formed from the retained austenite as the nuclei.
• martensite and bainite contain numerous dislocations, and B that has been dissolved therein diffuses rapidly into austenite grain boundaries through the dislocations during the second annealing; thus, uniform segregation of B occurs.
• the present invention has been made based on the aforementioned findings.
• the gist of features of the present invention is as follows.
• a high-strength steel sheet and a member having a tensile strength TS of 1180 MPa or more and excellent stretch flangeability and toughness, and methods for producing the same can be provided.
• a steel sheet according to an embodiment has a chemical composition containing, in mass%, C: 0.10% or more and 0.30% or less, Si: 0.20% or more and 1.20% or less, Mn: 2.5% or more and 4.0% or less, P: 0.050% or less, S: 0.020% or less, Al: 0.10% or less, N: 0.01% or less, Ti: 0.100% or less, Nb: 0.002% or more and 0.050% or less, B: 0.0015% or more and 0.0040% or less, and the balance being Fe and incidental impurities, in which the chemical composition satisfies formula (1) below, a total area fraction of martensite and bainite is 95% or more, a retained austenite area fraction is 5% or less, a ferrite area fraction is 1% or less, a prior-austenite grain size is 10 ⁇ m or less, a C concentration in a prior-austenite grain boundary is equal to or more than 1.5 times a C content in steel, a B concentration in
• [%N] and [%Ti] respectively represent a N content and a Ti content (mass%) in steel.
• C has an effect of strengthening the martensite-bainite microstructure.
• C segregates to the prior-austenite grain boundary by the reheating treatment after the second annealing step and has an effect of improving the toughness.
• TS tensile strength
• the C content is 0.10% or more.
• the C content is preferably 0.11% or more.
• carboborides of B and iron are formed during annealing, and a sufficient amount of B cannot segregate to the grain boundaries.
• the C content is 0.30% or less.
• the C content is preferably 0.28% or less.
• Si 0.20% or more and 1.20% or less
• Si is an element effective for solid solution strengthening and the Si content needs to be 0.20% or more.
• the Si content is 0.20% or more.
• the Si content is preferably 0.50% or more.
• the Si is a ferrite-stabilizing element, and ferrite is formed at a Si content exceeding 1.20%, thereby degrading strength, stretch flangeability, and toughness.
• the Si content is 1.20% or less.
• the Si content is preferably 1.10% or less.
• Mn 2.5% or more and 4.0% or less
• Mn is effective for improving hardenability.
• the Mn content is 2.5% or more.
• the Mn content is preferably 2.8% or more.
• segregation sites become excessively hard, and the toughness is degraded.
• the Mn content is 4.0% or less.
• the Mn content is preferably 3.5% or less.
• the P content is 0.050% or less.
• the P content is preferably 0.025% or less.
• the lower limit of the P content is not particularly specified and may be 0%; however, it involves a higher production cost to make the P content less than 0.001%, and thus the P content is preferably 0.001% or more.
• the S content is 0.020% or less.
• the S content is preferably 0.018% or less.
• the S content is more preferably 0.0040% or less and even more preferably 0.0020% or less.
• the lower limit of the S content is not particularly specified; however, it involves a higher production cost to make the S content less than 0.0001%, and thus the S content is preferably 0.0001% or more.
• Al is an element that acts as a deoxidizing agent, and, in order to obtain such an effect, the Al content is preferably 0.005% or more. Meanwhile, at an Al content exceeding 0.10%, ferrite tends to form and the strength decreases. Thus, the Al content is 0.10% or less. The Al content is preferably 0.05% or less.
• N forms nitrides with Nb and B and diminishes the effects of adding Nb and B.
• the N content is 0.01% or less.
• the N content is preferably 0.006% or less.
• the lower limit is not particularly specified; however, from the production cost perspective, the N content is preferably 0.0001% or more.
• Ti fixes N in the steel as TiN and suppresses formation of BN and NbN, and thus has an effect of enhancing the effects of adding Nb and B and improving the toughness and stretch flangeability.
• the Ti content is preferably 0.005% or more. Meanwhile, at a Ti content exceeding 0.100%, coarse Ti carbides form on the grain boundaries, and the toughness is degraded. Thus, the Ti content is 0.100% or less.
• the Ti content is preferably 0.050% or less.
• Nb 0.002% or more and 0.050% or less
• Nb dissolves or precipitates as fine carbides and suppresses growth of the austenite grains during annealing.
• Nb can make the crystal grain size finer, complicate the fracture paths, and improve the toughness.
• the Nb content is 0.002% or more.
• the Nb content is preferably 0.005% or more.
• the effects are saturated, coarse Nb carbides precipitate, and the toughness is degraded.
• the Nb content is 0.050% or less.
• the Nb content is preferably 0.040% or less.
• the B content is 0.0015% or more.
• the B content is preferably 0.0016% or more.
• carboborides are formed, and the toughness is degraded.
• the B content is 0.0040% or less.
• the B content is preferably 0.0030% or less.
• N which easily bonds with these elements, needs to be fixed by Ti.
• the molar fraction of N is to be smaller than the molar fraction of Ti.
• the N content and the Ti content in the steel are adjusted to satisfy formula (1) described above.
• ([%N]/14)/([%Ti]/47.9) is preferably 0.6 or less.
• [%N] and [%Ti] respectively represent the N content and the Ti content (mass%) in the steel.
• the balance other than the aforementioned components is Fe and incidental impurities. Note that, regarding the optional components described below, the effects of the present invention are not impaired as long as the contents thereof are lower than the lower limit values, and thus these optional elements contained in amounts below the lower limits are treated as incidental impurities.
• the chemical composition of the steel sheet according to this embodiment may further contain, in mass%, at least one element selected from V: 0.100% or less, Mo: 0.500% or less, Cr: 1.00% or less, Cu: 1.00% or less, Ni: 0.50% or less, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.200% or less, W: 0.400% or less, Zr: 0.0200% or less, Ca: 0.0200% or less, Mg: 0.0200% or less, Co: 0.020% or less, REM: 0.0200% or less, Te: 0.020% or less, Hf: 0.10% or less, and Bi: 0.200% or less.
• V forms fine carbides and has an effect of increasing the strength. At a V content exceeding 0.100%, coarse V carbides precipitate, and the toughness is degraded in some cases. Thus, when V is to be contained, the V content is 0.100% or less.
• the V content is preferably 0.080% or less and more preferably 0.060% or less.
• the lower limit of the V content is not particularly limited and may be 0.000%; however, since V has an effect of increasing the strength by forming fine carbides, the V content is preferably 0.001% or more.
• the V content is more preferably 0.005% or more and even more preferably 0.010% or more.
• Mo improves hardenability and has an effect of increasing the bainite and martensite area fractions. At a Mo content exceeding 0.500%, the effect saturates. Thus, when Mo is to be contained, the Mo content is 0.500% or less.
• the Mo content is preferably 0.200% or less and more preferably 0.150% or less.
• the lower limit of the Mo content is not particularly limited and may be 0.000%; however, since Mo has effects of improving the hardenability and increasing the bainite and martensite area fractions, the Mo content is preferably 0.010% or more.
• the Mo content is more preferably 0.020% or more and even more preferably 0.030% or more.
• Cr has effects of improving hardenability and increasing the bainite and martensite area fractions. At a Cr content exceeding 1.00%, the effects saturate. Thus, when Cr is to be contained, the Cr content is 1.00% or less.
• the Cr content is preferably 0.300% or less and more preferably 0.250% or less.
• the lower limit of the Cr content is not particularly limited and may be 0.000%; however, since Cr has effects of improving hardenability and increasing the bainite and martensite area fractions, the Cr content is preferably 0.01% or more.
• the Cr content is more preferably 0.015% or more and even more preferably 0.030% or more.
• Cu when dissolved, has an effect of increasing the strength. In addition, Cu has an effect of improving the delayed fracture resistance. At a Cu content exceeding 1.00%, grain boundary cracking is likely to occur. Thus, when Cu is to be contained, the Cu content is 1.00% or less.
• the Cu content is preferably 0.60% or less and more preferably 0.30% or less.
• the lower limit of the Cu content is not particularly limited and may be 0.000%; however, since Cu has an effect of increasing the strength when dissolved, the Cu content is preferably 0.01% or more.
• the Cu content is more preferably 0.02% or more and even more preferably 0.05% or more.
• Ni has an effect of improving the hardenability, but the effect saturates at a Ni content exceeding 0.50%.
• the Ni content is 0.50% or less.
• the Ni content is preferably 0.20% or less and more preferably 0.15% or less.
• the lower limit of the Ni content is not particularly limited and may be 0.00%; however, since Ni has an effect of increasing the hardenability, the Ni content is preferably 0.01% or more.
• the Ni content is more preferably 0.02% or more and even more preferably 0.03% or more.
• Sb has an effect of suppressing surface oxidation, nitridation, and decarburization of the steel sheet, but the effect saturates at a Sb content exceeding 0.200%.
• the Sb content is 0.200% or less.
• the Sb content is preferably 0.050% or less and more preferably 0.020% or less.
• the lower limit of the Sb content is not particularly limited and may be 0.000%; however, since Sb has an effect of suppressing surface oxidation, nitridation, and decarburization of the steel sheet, the Sb content is preferably 0.001% or more.
• the Sb content is more preferably 0.002% or more and even more preferably 0.005% or more.
• Sn has an effect of suppressing surface oxidation, nitridation, and decarburization of the steel sheet.
• the Sn content is 0.200% or less.
• the Sn content is preferably 0.050% or less and more preferably 0.020% or less.
• the lower limit of the Sn content is not particularly limited and may be 0.000%; however, since Sn has an effect of suppressing surface oxidation, nitridation, and decarburization of the steel sheet, the Sn content is preferably 0.001% or more.
• the Sn content is more preferably 0.002% or more and even more preferably 0.005% or more.
• Ta forms fine carbides and has an effect of increasing the strength. At a Ta content exceeding 0.200%, coarse Ta carbides precipitate, and the toughness is degraded in some cases. Thus, when Ta is to be contained, the Ta content is 0.200% or less.
• the Ta content is preferably 0.100% or less and more preferably 0.070% or less.
• the lower limit of the Ta content is not particularly limited and may be 0.000%; however, since Ta has an effect of increasing the strength by forming fine carbides, the Ta content is preferably 0.001% or more.
• the Ta content is more preferably 0.005% or more and even more preferably 0.010% or more.
• W forms fine carbides and has an effect of increasing the strength. At a W content exceeding 0.400%, coarse W carbides precipitate, and the toughness is degraded in some cases. Thus, when W is to be contained, the W content is 0.400% or less.
• the W content is preferably 0.300% or less and more preferably 0.250% or less.
• the lower limit of the W content is not particularly limited and may be 0.000%; however, since W has an effect of increasing the strength by forming fine carbides, the W content is preferably 0.001% or more.
• the W content is more preferably 0.005% or more and even more preferably 0.010% or more.
• the Zr has an effect of improving the toughness by spheroidizing inclusions and suppressing the stress concentration. At a Zr content exceeding 0.0200%, a large number of inclusions are formed, and the toughness is degraded in some cases. Thus, when Zr is to be contained, the Zr content is 0.0200% or less.
• the Zr content is preferably 0.0150% or less and more preferably 0.0100% or less.
• the lower limit of the Zr content is not particularly limited and may be 0.0000%; however, since Zr spheroidizes inclusions, suppresses stress concentration, and has an effect of improving the toughness, the Zr content is preferably 0.0001% or more.
• the Zr content is more preferably 0.0010% or more and even more preferably 0.0020% or more.
• Ca can be used as a deoxidizing material. At a Ca content exceeding 0.0200%, a large number of Ca-based inclusions are formed, and the toughness is degraded in some cases. Thus, when Ca is to be contained, the Ca content is 0.0200% or less.
• the Ca content is preferably 0.0100% or less and more preferably 0.0080% or less.
• the lower limit of the Ca content is not particularly limited and may be 0.0000%; however, since Ca can be used as a deoxidizing material, the Ca content is preferably 0.0001% or more.
• the Ca content is more preferably 0.0005% or more and even more preferably 0.0010% or more.
• Mg can be used as a deoxidizing material. At a Mg content exceeding 0.0200%, a large number of Mg-based inclusions are formed, and the toughness is degraded in some cases. Thus, when Mg is to be contained, the Mg content is 0.0200% or less.
• the Mg content is preferably 0.0100% or less and more preferably 0.0080% or less.
• the lower limit of the Mg content is not particularly limited and may be 0.0000%; however, since Mg can be used as a deoxidizing material, the Mg content is preferably 0.0001% or more.
• the Mg content is more preferably 0.0005% or more and even more preferably 0.0010% or more.
• Co has an effect of increasing the strength through solid solution strengthening. At a Co content exceeding 0.020%, the effect saturates. Thus, when Co is to be contained, the Co content is 0.020% or less.
• the Co content is preferably 0.015% or less and more preferably 0.010% or less.
• the lower limit of the Co content is not particularly limited and may be 0.000%; however, since Co has an effect of increasing the strength through solid solution strengthening, the Co content is preferably 0.001% or more.
• the Co content is more preferably 0.002% or more and even more preferably 0.005% or more.
• the REM has an effect of improving the toughness by spheroidizing inclusions and suppressing the stress concentration. At a REM content exceeding 0.0200%, a large number of inclusions are formed, and the toughness is degraded in some cases. Thus, when REM is to be contained, the REM content is 0.0200% or less.
• the REM content is preferably 0.0100% or less and more preferably 0.0050% or less.
• the lower limit of the REM content is not particularly limited and may be 0.0000%; however, since REM spheroidizes inclusions, suppresses stress concentration, and has an effect of improving the toughness, the REM content is preferably 0.0001% or more.
• the REM content is more preferably 0.0005% or more and even more preferably 0.0010% or more.
• REM refers to a lanthanoid element from scandium (Sc) having atomic number 21 to yttrium (Y) having atonic number 39 and lanthanum (La) having atomic number 57 to lutetium (Lu) having atomic number 71.
• a REM content is the total content of one or more of the elements selected from among the REMs described above.
• REM is preferably La, Ce, and Nd.
• Te has an effect of improving the toughness by spheroidizing inclusions and suppressing the stress concentration.
• Te content is 0.020% or less.
• the Te content is preferably 0.015% or less and more preferably 0.010% or less.
• the lower limit of the Te content is not particularly limited and may be 0.000%; however, since Te spheroidizes inclusions, suppresses stress concentration, and has an effect of improving the toughness, the Te content is preferably 0.001% or more.
• the Te content is more preferably 0.002% or more and even more preferably 0.004% or more.
• Hf has an effect of improving the toughness by spheroidizing inclusions and suppressing the stress concentration. At a Hf content exceeding 0.10%, a large number of inclusions are formed, and the toughness is degraded. Thus, when Hf is to be contained, the Hf content is 0.10% or less.
• the Hf content is preferably 0.08% or less and more preferably 0.05% or less.
• the lower limit of the Hf content is not particularly limited and may be 0.000%; however, since Hf spheroidizes inclusions, suppresses stress concentration, and has an effect of improving the toughness, the Hf content is preferably 0.01% or more.
• Bi has an effect of improving bendability by reducing segregation. At a Bi content exceeding 0.200%, a large number of inclusions are formed, and the bendability is degraded in some cases. Thus, when Bi is to be contained, the Bi content is 0.200% or less.
• the Bi content is preferably 0.100% or less and more preferably 0.050% or less.
• the Bi content is even more preferably 0.010% or less and even further more preferably 0.005% or less.
• the lower limit of the Bi content is not particularly limited and may be 0.000%; however, since Bi has effects of reducing segregation and improve bendability, the Bi content is preferably 0.001% or more.
• the Bi content is more preferably 0.002% or more and even more preferably 0.003% or more.
• Martensite and bainite are both hard phases and are necessary for achieving a TS of 1180 MPa or more.
• the total area fraction of martensite and bainite is 95% or more.
• the total area fraction of martensite and bainite is preferably 96% or more.
• the upper limit of the total area fraction of martensite and bainite is not particularly limited and may be 100%.
• Retained austenite may be included as the remaining microstructure other than martensite and bainite.
• the retained austenite area fraction is 5% or less.
• the retained austenite area fraction is preferably 4% or less.
• the retained austenite area fraction may be 0% or more than 0%.
• the ferrite area fraction is 1% or less.
• the ferrite area fraction may be 0%.
• the area fraction of each microstructure is measured as follows.
• the retained austenite area fraction is determined by taking a test specimen from each steel sheet, chemically polishing the rolled surface down to a sheet thickness t/4 position of the steel sheet, measuring the X-ray diffraction intensity and the diffraction peak positions of the polished surface by using an X-ray diffraction (XRD) instrument, and calculating a volume fraction therefrom as the retained austenite area fraction.
• XRD X-ray diffraction
• SEM images in three fields of view are taken from the observation surface at a magnification of 2000x with a field of view region of 57.1 ⁇ m ⁇ 42.9 ⁇ m.
• the total area fraction of martensite, bainite, and retained austenite, and the area fraction of the microstructure (ferrite) other than martensite, bainite, and retained austenite are determined by image analysis.
• the retained austenite area fraction determined by XRD is subtracted to determine the area fraction of martensite and bainite.
• the average of three fields of view is assumed to be the area fraction of the microstructure.
• Prior-austenite grain size 10 ⁇ m or less
• the prior-austenite grain size needs to be 10 ⁇ m or less.
• the prior-austenite grain size is 10 ⁇ m or less.
• the prior-austenite is preferably 9 ⁇ m or less.
• the lower limit of the prior-austenite grain size is not particularly limited; however, from the viewpoint of production technology, the prior-austenite grain size is preferably 1 ⁇ m or more.
• the prior-austenite grain size is more preferably 2 ⁇ m or more and even more preferably 3 ⁇ m or more.
• the grain size of the prior-austenite is measured as follows. A sheet thickness section parallel to the rolling direction of each steel sheet is polished and then etched with picral to prepare an observation surface. SEM images of the microstructure at the sheet thickness t/4 position in three fields of view are taken from the observation surface with SEM at a magnification of 2000x with a field of view region of 57.1 ⁇ m ⁇ 42.9 ⁇ m. From the obtained microstructure images, the grain size of each of the prior-austenite grains is determined by image analysis, and the average of the three fields of view is assumed to be the prior-austenite grain size (average crystal grain size).
• C concentration in prior-austenite grain boundary equal to or more than 1.5 times C content in steel
• C segregates to the prior-austenite grain boundaries and thereby strengthens the grain boundaries and improve the toughness.
• the aforementioned effect is obtained as long as the C concentration in the prior-austenite grain boundary is equal to or more than 1.5 times the C content in the steel.
• the C concentration in the prior-austenite grain boundary is equal to or more than 1.5 times the C content in the steel.
• the C concentration in the prior-austenite boundary satisfies the following equation.
• C concentration (mass%) in prior-austenite grain boundary/C content (mass%) in steel ⁇ 1.5.
• the C concentration in the prior-austenite grain boundary is preferably equal to or more than 2.0 times the C content in the steel and more preferably equal to or more than 2.5 times the C content in the steel.
• the upper limit of the C concentration in the prior-austenite grain boundary is not specified; however, in order to appropriately prevent precipitation of hard carbides and carboborides on the grain boundary and further improve the toughness, the C concentration is preferably less than 7% in mass%. More preferably, the C concentration is 2% or less in mass%.
• the B segregates to the prior-austenite grain boundaries and thereby strengthens the grain boundaries, and can improve the toughness.
• the aforementioned effect is obtained as long as the B concentration in the prior-austenite grain boundary is 0.05% or more in mass%.
• the B concentration in the prior-austenite grain boundary is 0.05% or more in mass%.
• the B concentration in the prior-austenite grain boundary is preferably 0.07% or more and more preferably 0.10% or more in mass%.
• the upper limit of the B concentration in the prior-austenite grain boundary is not specified; however, in order to appropriately prevent precipitation of hard carboborides on the grain boundary and further improve the toughness, the B concentration is preferably less than 6% in mass%. More preferably, the B concentration is 2% or less in mass%.
• the variation in B concentration within the same prior-austenite grain boundary is less than 0.010% in mass%.
• the variation is preferably 0.009% or less in mass%, and more preferably 0.008% or less in mass%. The smaller the variation, the better; however, from the viewpoint of the production technology, the variation may be 0.001% or more.
• the C concentration, the B concentration, and the variation in the prior-austenite grain boundary are measured as follows.
• a needle-shaped specimen is prepared from a prior-austenite grain boundary-containing region by SEM-FIB (focused ion beam) method.
• the obtained needle-shaped specimen is subjected to a 3DAP analysis by using a 3DAP instrument (LEAP 4000X Si produced by AMETEK).
• the measurement is carried out in laser mode.
• the specimen temperature is 80 K or lower.
• the number of C ions and the number of B ions detected from the prior-austenite grain boundary and the number of other ions are used to determine the C concentration and the B concentration in the prior-austenite grain boundary.
• the average value of two specimens is assumed to be the C concentration or the B concentration.
• circular regions not overlapping one another and having a diameter of 5 nm or more and 10 nm or less are set at five sites on a plane of the measured prior-austenite grain boundary, and, for each of the circles, a volume of a column extending in a direction normal to the grain boundary plane is set, the B concentration on the grain boundary is determined for each of the columns, and the standard deviation thereof is assumed to be the variation in B concentration.
• the prior-austenite grain size is very large with respect to a region sampled by the SEM-FIB method, the grain boundary targeted by one sampled specimen is from the same grain boundary in all cases. For example, whereas the prior-austenite grain size is about 9 ⁇ m, the diameter of the region of the sampled specimen is about 0.1 ⁇ m.
• the determined variation in B concentration and the like is the variation within the same grain boundary.
• a steel sheet having a tensile strength TS of 1180 MPa or more can be provided.
• the tensile strength TS of the steel sheet is preferably 1250 MPa or more.
• the aforementioned steel sheet may have a coating layer on at least one surface.
• the coating layer is preferably a hot-dip galvanized layer, a hot-dip galvannealed layer, or an electrogalvanized layer.
• the composition of the coating layer is not particularly limited and can be any known composition.
• the composition of the hot-dip galvanized layer is not particularly limited and may be a common composition.
• the coating layer has a composition that contains Fe: 20 mass% or less, Al: 0.001 mass% or more and 1.0 mass% or less, a total of 0 mass% or more and 3.5 mass% or less of at least one selected from the group consisting of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM, and the balance being Zn and incidental impurities.
• the Fe content in the coating layer is less than 7 mass%, and when the coating layer is a hot-dip galvannealed layer, in one example, the Fe content in the coating layer is 7 mass% or more and 15 mass% or less and more preferably 8 mass% or more and 13 mass% or less.
• the coating weight is not particularly limited and is preferably 20 to 80 g/m 2 per side of the steel sheet.
• the coating layer is formed on both surfaces, top and bottom surfaces, of a steel sheet (a high-strength cold rolled steel sheet).
• a method for producing a steel sheet includes: a hot rolling step of hot-rolling a steel slab having the chemical composition described above to obtain a hot rolled sheet; a pickling step of pickling the hot rolled sheet; a cold rolling step of cold-rolling the hot rolled sheet, which has been subjected to the pickling step, to obtain a cold rolled sheet; a first annealing step of heating the cold rolled sheet to a first heating temperature equal to or higher than an Ac3 point; a cooling step of starting cooling of the cold rolled sheet, which has been subjected to the first annealing step, from the first heating temperature and cooling the cold rolled sheet to a cooling stop temperature equal to or higher than 100°C and lower than an Ms temperature at an average cooling rate of 50°C/s or more; a first reheating step of performing heating to a first reheating temperature of 300°C or higher and 400°C or lower after the cooling step, holding the first reheating temperature for 60 s or longer, and performing cooling to room
• a steel slab having the aforementioned chemical composition is produced.
• a steel material is melted to prepare a molten steel having the aforementioned chemical composition.
• the melting method is not particularly limited, and any known melting method such as converter melting or electric arc melting is adoptable.
• the obtained molten steel is solidified to produce a steel slab (slab).
• the method for producing a steel slab from the molten steel is not particularly limited, and a method such as a continuous casting method, an ingot casting method, or a thin slab casting method can be employed.
• the steel slab may be temporarily cooled and reheated before hot rolling, or the casted steel slab may be continuously hot-rolled without being cooled to room temperature.
• the slab heating temperature is preferably 1100°C or higher and preferably 1300°C or lower in view of the rolling load and formation of scale.
• the slab heating method is not particularly limited, and, for example, the slab may be heated in a heating furnace according to a standard method.
• the aforementioned steel slab is hot-rolled into a hot rolled sheet.
• Hot-rolling is not particularly limited, and a common method may be employed. Cooling after the hot rolling is also not particularly limited and the hot rolled sheet is cooled to a coiling temperature. Next, the hot rolled sheet is coiled into a coil.
• the coiling temperature is preferably 400°C or higher. This is because the hot rolled sheet can be easily coiled without increasing the strength as long as the coiling temperature is 400°C or higher.
• the coiling temperature is more preferably 550°C or higher.
• the coiling temperature is preferably 750°C or lower.
• the hot rolled sheet may be heat-treated for softening.
• the hot rolled sheet is pickled in the pickling step.
• the scale on the hot rolled sheet coiled into a coil can be removed.
• the method for removing the scale is not particularly limited, but in order to completely remove the scale, pickling is preferably performed while uncoiling the hot rolled coil.
• the pickling method is not particularly limited and may be any standard method.
• the hot rolled sheet is cold-rolled into a cold rolled sheet.
• a hot rolled sheet from which scale has been removed is cleaned as appropriate and then cold-rolled into a cold rolled sheet.
• the cold rolling method is not particularly limited and may be any standard method.
• the cold rolled sheet is heated to a first heating temperature equal to or higher than the Ac3 point and is annealed in an austenite single phase region.
• a first heating temperature equal to or higher than the Ac3 point
• ferrite is formed. This ferrite has dislocations reduced by annealing and thus there is no dislocation that serves as a boron (B) diffusion path during the second annealing; thus, it becomes difficult to uniformly segregate boron.
• excessive heating increases the austenite grain size to larger than 10 ⁇ m.
• the first heating temperature is equal to or higher than the Ac3 point.
• the first heating temperature is preferably equal to or higher than Ac3 point + 10°C and more preferably equal to or higher than Ac3 point + 20°C.
• the austenite microstructure formed by the second annealing has the same crystal structure as the austenite microstructure formed by the first annealing; thus, the first heating temperature is preferably 980°C or lower to make the austenite grain size to 10 ⁇ m or less even in the first annealing.
• the first heating temperature is more preferably 950°C or lower.
• the Ac3 point is determined by the following equation.
• First reheating step heating to first reheating temperature of 300°C or higher and 400°C or lower, holding first reheating temperature for 60 s or longer, and performing cooling to room temperature
• the cold rolled sheet is subjected to a partial quenching-partitioning treatment in the cooling step.
• the cooling stop temperature of the partial quenching is lower than 100°C, martensite transformation occurs before C partitioning, and a sufficient amount of retained austenite cannot be obtained before the second annealing.
• the amount of the retained austenite is insufficient, austenite having a different orientation than that in the first annealing is formed, and it is difficult to have boron (B) uniformly segregate to the grain boundaries of such austenite.
• the cooling stop temperature in the cooling step is 100°C or higher.
• the cooling stop temperature is preferably 120°C or higher and more preferably 150°C or higher. Meanwhile, martensite transformation does not occur at a cooling stop temperature equal to or higher than the Ms temperature, and thus carbon partitioning does not occur during the subsequent reheating and retained austenite is not formed before the second annealing. When the amount of the retained austenite is insufficient, austenite having a different orientation than that in the first annealing is formed, and it is difficult to have boron (B) uniformly segregate to the grain boundaries of such austenite. Thus, the cooling stop temperature in the cooling step is lower than the Ms temperature.
• the cooling stop temperature is preferably equal to or lower than Ms temperature - 20°C and more preferably equal to or lower than Ms temperature - 30°C.
• the average cooling rate is 50°C/s or more.
• the average cooling rate is preferably 60°C/s or more and more preferably 70°C/s or more.
• the upper limit for the average cooling rate is not particularly limited, but the cooling stop temperature becomes difficult to control when the cooling rate is excessively high; thus, the average cooing rate is preferably 1000°C/s or less and more preferably 200°C/s or less.
• the average cooling rate (°C/s) in the cooling step is "(first heating temperature (°C)) - (cooling stop temperature (°C))/(cooling time (sec) from first heating temperature (°C) to cooling stop temperature (°C))".
• the first reheating temperature is 300°C or higher.
• the first reheating temperature is preferably 310°C or higher and more preferably 320°C or higher. Meanwhile, when the first reheating temperature exceeds 400°C, non-transformed austenite decomposes into cementite, and retained austenite is not formed before the second annealing.
• the first reheating temperature is 400°C or lower.
• the first reheating temperature is preferably 390°C or lower and more preferably 380°C or lower.
• the first reheating temperature holding time is 60 s or longer.
• the holding time is preferably 80 s or longer and more preferably 100 s or longer.
• the holding time at the first reheating temperature is preferably shorter than 900 s and more preferably 600 s or shorter.
• the room temperature that serves as the cooling stop temperature in the first reheating step is not particularly limited and can be 5 to 50°C.
• the steel sheet (cold rolled sheet) obtained as above is again heated to a second heating temperature equal to or higher than the Ac3 point and annealed.
• austenite having the same orientation as austenite formed in the first annealing is formed from the retained austenite as the nuclei.
• boron (B) that has dissolved through the dislocations of martensite before austenite transformation rapidly diffuses into the austenite grain boundaries, and the boron segregation becomes uniform.
• the second heating temperature is lower than the Ac3 point, ferrite is formed, and the strength decreases.
• the second heating temperature is equal to or higher than the Ac3 point.
• the second heating temperature is preferably equal to or higher than Ac3 point + 10°C and more preferably equal to or higher than Ac3 point + 20°C. Meanwhile, excessive heating increases the austenite grain size to over 10 ⁇ m and degrades the toughness; thus, the second heating temperature is preferably 980°C or lower. The second heating temperature is more preferably 950°C or lower.
• cooling to room temperature may be performed or an overaging treatment may be performed by using the facility in the annealing line in order to perform tempering after formation of a sufficient amount of martensite in the subsequent second reheating step described below.
• the steel sheet is cooled after the second annealing step, heated to a second reheating temperature of 70°C or higher and 200°C or lower, and held at the second reheating temperature for 600 s or longer.
• the cooling conditions before heating (reheating) here are not particularly limited; however, from the viewpoints of suppressing ferrite transformation and suppressing the strength degradation, cooling is preferably performed at an average cooling rate of 30°C/s or more. Furthermore, cooling is preferably performed to room temperature. Note that the average cooling rate (°C/s) in this step is "(second heating temperature (°C)) - (cooling stop temperature (°C))/(cooling time (sec) from second heating temperature (°C) to cooling stop temperature (°C))".
• the second reheating temperature After cooling, heating (reheating) is performed to the second reheating temperature. Reheating causes C as well as B to segregate to the prior ⁇ grain boundaries and improves the toughness.
• reheating temperature At a second reheating temperature lower than 70°C, diffusion of C is slow and segregation of C is insufficient.
• the second reheating temperature is 70°C or higher.
• the second reheating temperature is preferably 90°C or higher.
• the second reheating temperature is 200°C or lower.
• the second reheating temperature is preferably 190°C or lower.
• the holding time at the second reheating temperature (reheating holding time) is shorter than 600 s, diffusion of C is slow and segregation of C is insufficient.
• the holding time at the second reheating temperature is 600 s or longer.
• the holding time at the second reheating temperature is preferably 800 s or longer.
• the upper limit of the holding time at the second reheating temperature is not particularly limited; however, in order to prevent precipitation of carboborides, the second reheating temperature is preferably 43200 s or shorter (0.5 day or shorter).
• At least one of the surfaces of the steel sheet may be coated in a coating step to obtain a steel sheet (high-strength coated steel sheet).
• the steel sheet subjected to the coating step may be heat-treated to alloy the coating layer of the steel sheet and to obtain an alloyed coated steel sheet.
• the production conditions other than those described above may be any standard conditions.
• a steel sheet according to this embodiment obtained as described above preferably has a sheet thickness of 0.5 mm or more.
• the sheet thickness is preferably 2.0 mm or less.
• a member that uses the aforementioned steel sheet in at least some part can be provided.
• the aforementioned steel sheet may be formed into an intended shape by a pressing process into an automotive part.
• the automotive part may include, as a raw material, a steel sheet other than the steel sheet of this embodiment.
• a high-strength steel sheet having a TS of 1180 MPa or more and excellent stretch flangeability and toughness can be provided, a member that has a TS of 1180 MPa or more as well as excellent stretch flangeability and toughness can be provided.
• the steel sheet according to this embodiment is suitable for use in automotive parts that contribute to weight reduction of automotive bodies.
• the steel sheet according to this embodiment is particularly suitable for use in, among automotive parts, members that are used in frame structure parts or reinforcing parts in general.
• the aforementioned method for producing a member includes a step of subjecting the aforementioned steel sheet to at least one of a forming process or a joining process to obtain a member.
• 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.
• Annealing was performed under the conditions shown in Table 2 using a continuous hot-dip galvanizing line to obtain steel sheets (cold rolled steel sheets (CR), hot-dip galvanized steel sheets (GI), and hot-dip galvannealed steel sheets (GA)).
• the hot-dip galvanized steel sheets were dipped in a coating bath at 460°C, and the coating weight was 35 g/m 2 per side.
• Hot-dip galvannealed steel sheets were produced by adjusting the coating amount to 45 g/m 2 per side and then performing an alloying treatment of holding 520°C for 40 s. The obtained steel sheets except for the steel sheet No.
• the total area fraction of martensite and bainite, the retained austenite area fraction, the ferrite area fraction, the prior-austenite grain size, the C concentration in the prior-austenite grain boundary, the B concentration in the prior-austenite grain boundary, and the variation in B concentration within the same grain boundary of prior-austenite were evaluated according to the methods described above.
• the tensile strength TS, the stretch flangeability, and the toughness were evaluated by the methods described below. The results are shown in Table 3.
• the obtained steel sheets were subjected to a tensile test in conformity with JIS Z 2241 (2011).
• a JIS No. 5 test piece for tensile test having a longitudinal direction orienting in a direction orthogonal to the rolling direction was taken, and was subjected to a tensile test to measure the tensile strength (TS).
• TS tensile strength
• a tensile strength TS of 1180 MPa or more was determined to be good tensile strength.
• a Charpy impact test was carried out in conformity with JIS Z 2242 (2016). From each of the obtained steel sheets, a test specimen having a width of 10 mm, a length of 55 mm, and a 90° V notch having a notch depth of 2 mm and being disposed at the lengthwise center was taken such that the direction in which the V notch was formed was perpendicular to the rolling direction of the steel sheet. Subsequently, a Charpy impact test was carried out in a test temperature range of -120 to +120°C. A transition curve was determined from the obtained percent brittle fracture, and the temperature at a percent brittle fracture of 50% was determined to be the brittle-to-ductile transition temperature.
• the cases where the brittle-to-ductile transition temperature obtained by the Charpy test was -40°C or lower were evaluated as having good toughness.
• the cases where the brittle-to-ductile transition temperature was -40°C or lower are indicated as "excellent” in toughness, and the cases where the brittle-to-ductile transition temperature exceeded -40°C are indicated as “poor” in toughness.
• a 100 mm W ⁇ 100 m L test specimen was taken from the steel sheet (cold rolled steel sheet or coated steel sheet) and a hole expansion test was conducted in conformity with JIS Z 2256 (2010).
• a 10 mm ⁇ hole was punched out in a test specimen under the condition of a clearance of 12 ⁇ 1%, a conical punch having a 60° apex angle was elevated to expand hole and stopped as soon as cracks occurred in the sheet thickness direction, and the hole expansion ratio ⁇ was determined from the equation below from the hole diameter after cracking and the hole diameter before the test.
• Table 3 indicates that the examples of the present invention have a tensile strength TS of 1180 MPa or more and excellent stretch flangeability and toughness. In contrast, in the comparative examples, at least one of the tensile strength TS, the stretch flangeability, and the toughness was poor.
• 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 and excellent stretch flangeability and toughness comparable to those of the steel sheet of the examples of the present invention since the steel sheet of the examples of the present invention had high strength and excellent stretch flangeability and toughness.

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