CN116829752A - High-strength steel sheet and method for producing same - Google Patents
High-strength steel sheet and method for producing same Download PDFInfo
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- CN116829752A CN116829752A CN202180093135.5A CN202180093135A CN116829752A CN 116829752 A CN116829752 A CN 116829752A CN 202180093135 A CN202180093135 A CN 202180093135A CN 116829752 A CN116829752 A CN 116829752A
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 116
- 239000010959 steel Substances 0.000 title claims abstract description 116
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 21
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 118
- 230000000717 retained effect Effects 0.000 claims abstract description 96
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 75
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 40
- 238000007747 plating Methods 0.000 claims abstract description 37
- 239000000203 mixture Substances 0.000 claims abstract description 28
- 238000011282 treatment Methods 0.000 claims abstract description 28
- 229910001563 bainite Inorganic materials 0.000 claims abstract description 18
- 238000001816 cooling Methods 0.000 claims description 38
- 230000009466 transformation Effects 0.000 claims description 38
- 238000003303 reheating Methods 0.000 claims description 28
- 238000005096 rolling process Methods 0.000 claims description 27
- HCHKCACWOHOZIP-UHFFFAOYSA-N Zinc Chemical compound [Zn] HCHKCACWOHOZIP-UHFFFAOYSA-N 0.000 claims description 20
- 229910052725 zinc Inorganic materials 0.000 claims description 20
- 239000011701 zinc Substances 0.000 claims description 20
- 238000005275 alloying Methods 0.000 claims description 12
- 229910052720 vanadium Inorganic materials 0.000 claims description 10
- 238000005097 cold rolling Methods 0.000 claims description 9
- 230000007704 transition Effects 0.000 claims description 9
- 229910052721 tungsten Inorganic materials 0.000 claims description 5
- 239000012535 impurity Substances 0.000 claims description 4
- 229910052757 nitrogen Inorganic materials 0.000 claims description 3
- 229910052698 phosphorus Inorganic materials 0.000 claims description 2
- 238000000137 annealing Methods 0.000 description 19
- 230000000694 effects Effects 0.000 description 14
- 238000010438 heat treatment Methods 0.000 description 14
- 230000007423 decrease Effects 0.000 description 13
- 239000010410 layer Substances 0.000 description 13
- 238000000034 method Methods 0.000 description 13
- 238000005452 bending Methods 0.000 description 11
- 238000005098 hot rolling Methods 0.000 description 11
- 229910052799 carbon Inorganic materials 0.000 description 10
- 238000012360 testing method Methods 0.000 description 10
- 239000011248 coating agent Substances 0.000 description 8
- 238000000576 coating method Methods 0.000 description 8
- 230000007547 defect Effects 0.000 description 8
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 8
- 239000002244 precipitate Substances 0.000 description 8
- 230000009467 reduction Effects 0.000 description 8
- 230000002829 reductive effect Effects 0.000 description 8
- 238000005728 strengthening Methods 0.000 description 8
- 238000005246 galvanizing Methods 0.000 description 7
- 239000000463 material Substances 0.000 description 7
- 239000013078 crystal Substances 0.000 description 6
- 230000006872 improvement Effects 0.000 description 6
- 229910052748 manganese Inorganic materials 0.000 description 6
- 238000005554 pickling Methods 0.000 description 5
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 4
- 239000002253 acid Substances 0.000 description 4
- 239000000470 constituent Substances 0.000 description 4
- 230000000670 limiting effect Effects 0.000 description 4
- 229910052750 molybdenum Inorganic materials 0.000 description 4
- 229910052758 niobium Inorganic materials 0.000 description 4
- 238000005406 washing Methods 0.000 description 4
- 230000002411 adverse Effects 0.000 description 3
- 229910045601 alloy Inorganic materials 0.000 description 3
- 239000000956 alloy Substances 0.000 description 3
- 238000005266 casting Methods 0.000 description 3
- 229910052804 chromium Inorganic materials 0.000 description 3
- 238000009749 continuous casting Methods 0.000 description 3
- 238000005336 cracking Methods 0.000 description 3
- 230000006866 deterioration Effects 0.000 description 3
- 239000000446 fuel Substances 0.000 description 3
- 238000000227 grinding Methods 0.000 description 3
- 230000000977 initiatory effect Effects 0.000 description 3
- 230000003647 oxidation Effects 0.000 description 3
- 238000007254 oxidation reaction Methods 0.000 description 3
- 229910001562 pearlite Inorganic materials 0.000 description 3
- 238000001556 precipitation Methods 0.000 description 3
- 239000000523 sample Substances 0.000 description 3
- 239000006104 solid solution Substances 0.000 description 3
- 230000000087 stabilizing effect Effects 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- LFQSCWFLJHTTHZ-UHFFFAOYSA-N Ethanol Chemical compound CCO LFQSCWFLJHTTHZ-UHFFFAOYSA-N 0.000 description 2
- 229910001335 Galvanized steel Inorganic materials 0.000 description 2
- 238000002441 X-ray diffraction Methods 0.000 description 2
- 230000032683 aging Effects 0.000 description 2
- 229910052787 antimony Inorganic materials 0.000 description 2
- 229910001567 cementite Inorganic materials 0.000 description 2
- 238000006243 chemical reaction Methods 0.000 description 2
- 239000010960 cold rolled steel Substances 0.000 description 2
- 238000007796 conventional method Methods 0.000 description 2
- 229910052802 copper Inorganic materials 0.000 description 2
- 238000009792 diffusion process Methods 0.000 description 2
- 239000012467 final product Substances 0.000 description 2
- 239000008397 galvanized steel Substances 0.000 description 2
- 230000014509 gene expression Effects 0.000 description 2
- 238000000265 homogenisation Methods 0.000 description 2
- 230000001771 impaired effect Effects 0.000 description 2
- 229910052742 iron Inorganic materials 0.000 description 2
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 2
- 238000005461 lubrication Methods 0.000 description 2
- 229910052759 nickel Inorganic materials 0.000 description 2
- 238000005121 nitriding Methods 0.000 description 2
- 239000011148 porous material Substances 0.000 description 2
- 238000003825 pressing Methods 0.000 description 2
- 238000012545 processing Methods 0.000 description 2
- 239000000047 product Substances 0.000 description 2
- 238000010791 quenching Methods 0.000 description 2
- 230000000171 quenching effect Effects 0.000 description 2
- 239000002344 surface layer Substances 0.000 description 2
- 229910052718 tin Inorganic materials 0.000 description 2
- 239000013585 weight reducing agent Substances 0.000 description 2
- 241000219307 Atriplex rosea Species 0.000 description 1
- GRYLNZFGIOXLOG-UHFFFAOYSA-N Nitric acid Chemical compound O[N+]([O-])=O GRYLNZFGIOXLOG-UHFFFAOYSA-N 0.000 description 1
- 229910001035 Soft ferrite Inorganic materials 0.000 description 1
- 230000002159 abnormal effect Effects 0.000 description 1
- 238000004458 analytical method Methods 0.000 description 1
- 238000005279 austempering Methods 0.000 description 1
- 230000003749 cleanliness Effects 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 239000002131 composite material Substances 0.000 description 1
- 238000012937 correction Methods 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 238000000354 decomposition reaction Methods 0.000 description 1
- 230000003247 decreasing effect Effects 0.000 description 1
- 230000003111 delayed effect Effects 0.000 description 1
- 230000002542 deteriorative effect Effects 0.000 description 1
- 238000011161 development Methods 0.000 description 1
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- 238000010191 image analysis Methods 0.000 description 1
- 238000007654 immersion Methods 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 229910017604 nitric acid Inorganic materials 0.000 description 1
- 238000005498 polishing Methods 0.000 description 1
- 229910001568 polygonal ferrite Inorganic materials 0.000 description 1
- 230000008569 process Effects 0.000 description 1
- 238000001953 recrystallisation Methods 0.000 description 1
- 239000011347 resin Substances 0.000 description 1
- 229920005989 resin Polymers 0.000 description 1
- 229920006395 saturated elastomer Polymers 0.000 description 1
- 238000001878 scanning electron micrograph Methods 0.000 description 1
- 238000005204 segregation Methods 0.000 description 1
- 239000007787 solid Substances 0.000 description 1
- 239000000243 solution Substances 0.000 description 1
- 239000007921 spray Substances 0.000 description 1
- 238000005482 strain hardening Methods 0.000 description 1
- 230000003746 surface roughness Effects 0.000 description 1
- 238000009864 tensile test Methods 0.000 description 1
- 150000003568 thioethers Chemical class 0.000 description 1
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 1
Classifications
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—Hardening; Quenching with or without subsequent tempering
- C21D1/185—Hardening; Quenching with or without subsequent tempering from an intercritical temperature
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—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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- 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
- C21D1/78—Combined heat-treatments not provided for above
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- 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
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—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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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—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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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0273—Final recrystallisation annealing
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/002—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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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- 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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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—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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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/60—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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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
- C23C2/06—Zinc or cadmium or alloys based thereon
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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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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
Abstract
The purpose of the present invention is to provide a high-strength steel sheet which has a TS of 980MPa or more, has excellent ductility, hole expansibility and bendability, and does not reduce ductility after plating treatment, and a method for producing the same. A high-strength steel sheet having a predetermined composition, a steel structure in which ferrite is 1% or more and 40% or less, fresh martensite is 1% or more and 20% or less, the sum of bainite and tempered martensite is 35% or more and 90% or less, and retained austenite is 6% or more, the value obtained by dividing the average Mn content (mass%) in retained austenite by the average Mn content (mass%) in ferrite is 1.1 or more, the value obtained by dividing the average C content (mass%) in retained austenite having an aspect ratio of 2.0 or more by the average C content (mass%) in ferrite is 3.0 or more, and the C content in the whole retained austenite is divided by the C content (mass%) in the whole retained austeniteT 0 The amount of C in the composition is 1.0 or more.
Description
Technical Field
The present invention relates to a high-strength steel sheet excellent in formability suitable for use as a member in the industrial fields such as automobiles and electric appliances, and a method for producing the same, and particularly, it is desired to obtain a high-strength steel sheet excellent in ductility and hole expansibility and bendability, as well as having a TS (tensile strength) of 980MPa or more.
Background
In recent years, from the viewpoint of global environment protection, improvement of fuel efficiency of automobiles has become an important issue. Therefore, the trend of reducing the weight of the vehicle body itself is becoming more and more active in achieving a reduction in the thickness of the vehicle body material due to a higher strength. On the other hand, since the strength of a steel sheet is increased to reduce the formability, development of a material having both high strength and high formability is desired.
As a steel sheet having high strength and excellent high ductility, a high strength steel sheet in which transformation is induced by processing of retained austenite has been proposed. Such a steel sheet has a structure having retained austenite, and is easy to form by retained austenite during forming of the steel sheet, and has high strength because the retained austenite is martensitic after forming.
For example, patent document 1 proposes a high-strength steel sheet having a tensile strength of 1000MPa or more and a total Elongation (EL) of 30% or more, which has very high ductility by utilizing a work-induced transformation of retained austenite. Such steel sheets are produced by austenitizing a steel sheet containing C, si, and Mn as essential components, then quenching the steel sheet in a bainite transformation temperature range, and holding the steel sheet isothermally, so-called austempering. The retained austenite is produced by enriching C into austenite by the isothermal quenching treatment, but in order to obtain a large amount of retained austenite, a large amount of C exceeding 0.3% needs to be added. However, when the C concentration in steel increases, the spot weldability decreases, and particularly when the C concentration exceeds 0.3%, the decrease is remarkable, and it is difficult to put the steel sheet to practical use as a steel sheet for automobiles. In addition, the above patent documents do not consider hole expansibility because the main purpose of the above patent documents is to improve the ductility of a high-strength steel sheet.
In patent document 2, a steel containing Mn of 4 wt% or more and 6 wt% or less is used, and heat treatment is performed in a duplex region of ferrite and austenite, whereby a high strength-ductility balance is obtained. However, patent document 2 does not address the improvement of ductility due to the enrichment of Mn into non-phase-transformed austenite, and there is room for improvement in workability.
Patent document 3 discloses that: the heat treatment in the ferrite-austenite dual-phase region is performed using a steel containing 3.0 mass% or more and 7.0 mass% or less of Mn. As a result, mn is enriched in the non-phase-transformed austenite to form stable retained austenite, thereby improving the total elongation. However, since the heat treatment time is short and the diffusion rate of Mn is slow, it is presumed that the enrichment of Mn is insufficient in order to have both hole expansibility and bendability in addition to the elongation.
Further, patent document 4 discloses: the hot rolled sheet is subjected to a long-time heat treatment in a ferrite-austenite dual-phase region using a steel containing 0.50 mass% or more and 12.00 mass% or less of Mn. As a result, retained austenite having a large aspect ratio is formed, which promotes enrichment of Mn into non-phase-transformed austenite, and uniform elongation is improved. However, there has been no study on improvement of hole expansibility, and on both of flexibility and elongation. In the plating step and the alloying treatment step, austenite is easily decomposed, and therefore, it is difficult to secure a desired amount of retained austenite.
Prior art literature
Patent literature
Patent document 1: japanese patent laid-open No. 61-157625
Patent document 2: japanese patent laid-open No. 1-259120
Patent document 3: japanese patent laid-open No. 2003-138345
Patent document 4: japanese patent No. 6123966
Disclosure of Invention
Problems to be solved by the invention
The present invention has been made in view of the above-described circumstances, and an object thereof is to provide a high-strength steel sheet having a TS (tensile strength) of 980MPa or more, excellent formability, and no reduction in ductility after plating treatment, and a method for producing the same. The formability described herein means ductility, hole expansibility and bendability.
Means for solving the problems
The present inventors have conducted intensive studies to solve the above problems and as a result, have found the following matters in view of the composition of the steel sheet and the manufacturing method in order to manufacture a high-strength steel sheet having excellent formability.
That is, it has been found that the composition of the components containing Mn in an amount of 2.00 mass% or more and 8.00 mass% or less and other alloying elements such as Ti is suitably adjusted, and after hot rolling, ac is added as required 1 The cold rolling is performed by performing pickling treatment if necessary while maintaining the temperature range below the transformation point for more than 1800 seconds. Then, at Ac 3 The temperature of the transformation point is maintained in a range of-50 ℃ or more for 20s or more and 1800s or less, and then cooled to a cooling stop temperature of not more than the martensitic transformation start temperature, and then reheated to a reheating temperature in a range of 120 ℃ or more and 450 ℃ or less. Then, it is important to form C-enriched membranous austenite as a core of fine residual austenite having a large aspect ratio and significantly enriched in Mn and C in the subsequent annealing step by maintaining the reheating temperature for 2s or more and 1800s or less and then cooling the obtained product to room temperature.
After the cooling, the mixture was cooled by Ac 1 The temperature of the transformation point is maintained in a range of-20 ℃ or more for 20s or more and 600s or less, and then cooled to a cooling stop temperature of not more than the martensitic transformation start temperature, and then reheated to a reheating temperature in a range of 120 ℃ or more and 480 ℃ or less. Then, the mixture is kept at the reheating temperature for 2 to 600 seconds, and then cooled to room temperature. As a result, it was found that a high-strength steel sheet having excellent formability, characterized in that ferrite of 1% or more and 40% or less, fresh martensite of 1% or more and 20% or less, and Belleville can be obtained in terms of area ratio A steel structure in which the sum of the tempered martensite and the bulk is 35% or more and 90% or less, and the retained austenite is 6% or more, wherein the value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite is 1.1 or more, the value obtained by dividing the average C amount (mass%) in the retained austenite having an aspect ratio of 2.0 or more by the average C amount (mass%) in the ferrite is 3.0 or more, and the C amount in the entire retained austenite is divided by T 0 The amount of C in the composition is 1.0 or more.
The present invention has been completed based on the above-described findings, and the gist thereof is as follows.
[1] A high strength steel sheet, comprising:
contains C in mass%: 0.030% or more and 0.250% or less, si:0.01% or more and 3.00% or less, mn:2.00% or more and 8.00% or less, P:0.100% or less, S: below 0.0200%, N: less than 0.0100%, al:0.001% or more and 2.000% or less, the balance being Fe and unavoidable impurities; and a steel structure in which ferrite is 1% or more and 40% or less, fresh martensite is 1% or more and 20% or less, the sum of bainite and tempered martensite is 35% or more and 90% or less, and retained austenite is 6% or more, wherein the value obtained by dividing the average Mn amount (mass%) in retained austenite by the average Mn amount (mass%) in ferrite is 1.1 or more, the value obtained by dividing the average C amount (mass%) in retained austenite having an aspect ratio of 2.0 or more by the average C amount (mass%) in ferrite is 3.0 or more, and the C amount in the whole retained austenite is divided by T 0 The amount of C in the composition is 1.0 or more.
[2] The high-strength steel sheet according to [1], wherein the composition of the above components further contains, in mass%, a composition selected from the group consisting of Ti: less than 0.200%, nb: less than 0.200%, V: less than 0.500%, W:0.500% or less, B: less than 0.0050%, ni: less than 1.000%, cr: less than 1.000%, mo: less than 1.000%, cu: less than 1.000%, sn:0.200% or less, sb: less than 0.200%, ta: less than 0.100%, zr: less than 0.200%, ca: less than 0.0050%, mg: less than 0.0050%, REM:0.0050% or less of at least one element.
[3] The high-strength steel sheet according to [1] or [2], wherein the value obtained by dividing the area ratio of bulk retained austenite by the area ratio of all retained austenite and bulk fresh martensite is 0.5 or less.
[4] The high-strength steel sheet according to [1] to [3], which further comprises a zinc-plated layer on the surface.
[5] The high-strength steel sheet according to [4], wherein the zinc plating layer is an alloyed zinc plating layer.
[6]A method for producing a high-strength steel sheet, which comprises the steps of [1]]～[3]The method for producing a high-strength steel sheet according to any one of claims, wherein the steel sheet has a composition of [1]]Or [2]]The steel billet composed of the above components is heated, hot rolled under the condition that the temperature of the finish rolling outlet side is more than 750 ℃ and less than 1000 ℃, coiled and cold rolled under the condition that the temperature is more than 300 ℃ and less than 750 ℃, and then Ac is carried out 3 Maintaining the temperature of transformation point at-50deg.C for 20s or more and 1800s or less, cooling to cooling stop temperature of martensite transformation start temperature or less, reheating to reheating temperature of 120deg.C or more and 450 deg.C or less, maintaining the reheating temperature for 2s or more and 1800s or less, cooling to room temperature, and cooling to Ac 1 The temperature range of the transformation point-20 ℃ is maintained at 20-600 s, cooled to a cooling stop temperature of not higher than the martensitic transformation start temperature, reheated to a reheating temperature in the range of 120-480 ℃, maintained at the reheating temperature for 2-600 s, and cooled to room temperature.
[7] The method for producing a high-strength steel sheet according to [6], wherein a galvanization treatment is further performed.
[8] The method of producing a high-strength steel sheet according to [7], wherein the alloying treatment is performed at 450℃or more and 600℃or less after the galvanization treatment.
[9]According to [6]]～[8]The method for producing a high-strength steel sheet according to any one of the above, wherein Ac is a value after coiling and before cold rolling 1 The temperature range below the phase transition point remains above 1800s.
Effects of the invention
According to the present invention, a high-strength steel sheet having a TS (tensile strength) of 980MPa or more, formability after plating, particularly excellent ductility, hole expansibility and bendability, and further having no reduction in ductility after plating can be obtained. By applying the high-strength steel sheet obtained by the production method of the present invention to, for example, an automobile structural member, improvement in fuel efficiency due to weight reduction of the automobile body can be achieved, and industrial utilization value is extremely high.
Detailed Description
The present invention will be specifically described below. The "%" indicating the content of the constituent elements refers to "% by mass" unless otherwise specified.
(1) The reasons why the composition of the steel is limited to the above range in the present invention will be described.
C:0.030% or more and 0.250% or less
C is an element necessary for increasing strength by generating a low-temperature transformation phase such as martensite. In addition, C is an element effective for improving the stability of retained austenite and improving the ductility of steel. If the C content is less than 0.030%, ferrite is excessively generated, and thus the desired strength cannot be obtained. In addition, it is difficult to secure a sufficient area ratio of retained austenite, and good ductility is not obtained. On the other hand, when C exceeds 0.250% and is excessively contained, the area ratio of hard martensite becomes excessively large, and in the hole expansion test, the micro voids at the grain boundaries of martensite increase, and propagation of cracks proceeds, and hole expansibility decreases. Further, hardening of the welded portion and the heat affected portion is remarkable, and mechanical properties of the welded portion are lowered, so that spot weldability, arc weldability, and the like are deteriorated. From such a viewpoint, the C amount is set to 0.030% or more and 0.250% or less. The lower limit value is preferably 0.080% or more. The upper limit is preferably 0.200% or less.
Si:0.01% to 3.00%
Si improves work hardening ability of ferrite, and is therefore effective to ensure good ductility. When the Si content is less than 0.01%, the effect of addition is insufficient, and therefore the lower limit is set to 0.01%. However, excessive addition of Si exceeding 3.00% causes deterioration of ductility and bendability due to embrittlement of steel, and also causes deterioration of surface properties due to generation of red scale or the like. Further resulting in a decrease in the quality of the coating. Therefore, si is set to 0.01% or more and 3.00% or less. The lower limit is preferably 0.20% or more. The upper limit is preferably 2.00% or less, more preferably less than 1.20%.
Mn:2.00% to 8.00%
Mn is an extremely important element in the present invention. Mn is an element that stabilizes retained austenite, is effective in ensuring good ductility, and is an element that increases the strength of steel by solid solution strengthening. Such an effect was confirmed when the Mn content of the steel was 2.00% or more. However, excessive addition of Mn in an amount exceeding 8.00% deteriorates chemical conversion treatability and plating quality. From such a viewpoint, the Mn amount is set to 2.00% or more and 8.00% or less. The lower limit is preferably 2.30% or more, more preferably 2.50% or more. The upper limit is preferably 6.00% or less, more preferably 4.20% or less.
P: less than 0.100%
P is an element having a solid solution strengthening effect and capable of being added according to a desired strength. When the P content exceeds 0.100%, the weldability is deteriorated, and the alloying rate is lowered when the galvanized layer is subjected to the alloying treatment, thereby deteriorating the quality of the galvanized layer. The lower limit may be 0%, but is preferably 0.001% or more in terms of production cost. Therefore, the P amount is set to 0.100% or less. The lower limit is more preferably 0.005% or more. The preferable upper limit value is set to 0.050% or less.
S: less than 0.0200%
S segregates to grain boundaries to embrittle steel during hot working, and exists as sulfides to reduce local deformability. Therefore, the amount thereof needs to be set to 0.0200% or less, preferably 0.0100% or less, more preferably 0.0050% or less. The lower limit may be 0%, but is preferably 0.0001% or more in terms of production cost.
N:0.0100% or less
N is an element that deteriorates the aging resistance of steel. In particular, when the N content exceeds 0.0100%, deterioration of aging resistance becomes remarkable. The lower limit value may be 0% as the amount thereof is smaller, but the amount of N is preferably 0.0005% or more from the viewpoint of production cost. Therefore, the N amount is set to 0.0100% or less. More preferably, the content is set to 0.0010% or more. The upper limit of the N amount is preferably set to 0.0070% or less.
Al:0.001% or more and 2.000% or less
Al is an element effective for expanding a ferrite-austenite dual-phase region and reducing the annealing temperature dependency on mechanical properties, that is, for improving the material stability. When the content of Al is less than 0.001%, the effect of addition is insufficient, and therefore, the lower limit is set to 0.001%. Al is an element that functions as a deoxidizer and is effective for the cleanliness of steel, and is preferably added in the deoxidizing step. However, the addition of a large amount exceeding 2.000% increases the risk of cracking of the steel sheet during continuous casting, and decreases manufacturability. From such a viewpoint, the Al amount is set to 0.001% or more and 2.000% or less. The lower limit is preferably 0.025% or more, more preferably 0.200% or more. The upper limit is preferably 1.200% or less.
In addition, the above-mentioned components may contain, in mass%, a component selected from the group consisting of Ti: less than 0.200%, nb: less than 0.200%, V: less than 0.500%, W:0.500% or less, B: less than 0.0050%, ni: less than 1.000%, cr: less than 1.000%, mo: less than 1.000%, cu: less than 1.000%, sn:0.200% or less, sb: less than 0.200%, ta: less than 0.1000%, zr: less than 0.200%, ca: less than 0.0050%, mg: less than 0.0050%, REM:0.0050% or less of at least one element.
Ti: less than 0.200%
Ti is effective for precipitation strengthening of steel, and by increasing the strength of ferrite, the difference in hardness from the hard second phase (martensite or retained austenite) can be reduced, and further excellent hole expansibility can be ensured, so that it can be contained as needed. However, if the area ratio of the hard martensite exceeds 0.200%, the micro voids at the grain boundaries of the martensite increase and the propagation of cracks proceeds at the time of the hole expansion test, and the hole expansion property may be lowered. Therefore, when Ti is added, the addition amount is set to 0.200% or less. The lower limit is preferably 0.005% or more, more preferably 0.010% or more. The upper limit value is preferably set to 0.100% or less.
Nb: less than 0.200%, V: less than 0.500%, W: less than 0.500%
Nb, V, and W are effective for precipitation strengthening of steel, and as with the effect of Ti addition, by increasing the strength of ferrite, the difference in hardness from the hard second phase (martensite or retained austenite) can be reduced, and further excellent hole expansibility can be ensured, and therefore, can be contained as needed. However, if Nb exceeds 0.200% and V, W exceeds 0.500%, the area ratio of hard martensite becomes excessively large, and in the hole expansion test, micro voids at the grain boundaries of the martensite increase, propagation of cracks proceeds, and hole expansion properties may be reduced. Therefore, in the case of adding Nb, the addition amount thereof is set to 0.200% or less. The lower limit value of Nb is preferably set to 0.005% or more, more preferably 0.010% or more. The upper limit of Nb is preferably set to 0.100% or less. In the case of adding V, W, the addition amounts thereof were each set to 0.500% or less. The lower limit value of V, W is preferably set to 0.005% or more, more preferably 0.010% or more. The upper limit value of V, W is preferably set to 0.300% or less.
B: less than 0.0050%
B has an effect of suppressing the generation and growth of ferrite from austenite grain boundaries, and by increasing the strength of ferrite, the hardness difference with the hard second phase (martensite or retained austenite) can be reduced, and further excellent hole expansibility can be ensured, so that it can be contained as needed. However, when the content exceeds 0.0050%, the formability may be lowered. Therefore, when B is added, the addition amount is set to 0.0050% or less. The lower limit value is preferably set to 0.0003% or more, more preferably 0.0005% or more. The upper limit value is preferably set to 0.0030% or less.
Ni: less than 1.000 percent
Ni is an element for stabilizing retained austenite, is effective for securing better ductility, and is an element for further increasing the strength of steel by solid solution strengthening, and therefore may be contained as needed. On the other hand, when the addition amount exceeds 1.000%, the area ratio of the hard martensite becomes excessively large, and in the hole expansion test, the micro voids at the grain boundaries of the martensite increase, and the propagation of cracks proceeds, and the hole expansion property decreases. Therefore, in the case of adding Ni, the addition amount is set to 1.000% or less, preferably to 0.005% or more and 1.000% or less.
Cr: less than 1.000%, mo: less than 1.000 percent
Cr and Mo have an effect of improving the balance between strength and ductility, and therefore may be added as needed. However, cr and Mo each exceeds Cr:1.000%, mo: if 1.000% is excessively added, the area ratio of the hard martensite becomes excessively large, and in the hole expansion test, the micro voids at the grain boundaries of the martensite increase, and propagation of cracks proceeds, and hole expansion properties sometimes decrease. Therefore, in the case of adding these elements, the amounts thereof are each set to Cr: less than 1.000%, mo: less than 1.000%, preferably set to Cr:0.005% or more and 1.000% or less, mo:0.005% or more and 1.000% or less.
Cu: less than 1.000 percent
Cu is an element effective for strengthening steel, and can be used for strengthening steel as required within the range specified in the present invention. On the other hand, when the addition amount exceeds 1.000%, the area ratio of the hard martensite becomes excessively large, and in the hole expansion test, the micro voids at the grain boundaries of the martensite increase, and the propagation of cracks proceeds, and the hole expansion property decreases. Therefore, when Cu is added, the amount is set to 1.000% or less, preferably 0.005% or more and 1.000% or less.
Sn:0.200% or less, sb: less than 0.200%
Sn and Sb are added as needed from the viewpoint of suppressing decarburization of a region of about several tens μm of the steel sheet surface layer due to nitriding and oxidation of the steel sheet surface. Such nitriding and oxidation are suppressed, and the area ratio of martensite in the steel sheet surface is prevented from decreasing, and it is effective to ensure strength and material stability, and therefore, may be contained as needed. On the other hand, if any of these elements exceeds 0.200% and is excessively added, the toughness is lowered. Therefore, when Sn and Sb are added, the content thereof is set to 0.200% or less, preferably 0.002% or more and 0.200% or less, respectively.
Ta: less than 0.100%
Like Ti and Nb, ta forms alloy carbide and alloy carbonitride, contributing to higher strength. In addition, ta is considered to have the following effects: a part of the precipitate is solid-dissolved in Nb carbide or Nb carbonitride to form a composite precipitate such as (Nb, ta) (C, N), thereby remarkably suppressing coarsening of the precipitate and stabilizing the contribution of precipitation strengthening to strength. Therefore, ta may be contained as needed. On the other hand, even if Ta is excessively added, the effect of stabilizing the precipitate is saturated, and the alloy cost increases. Therefore, when Ta is added, the content is set to 0.100% or less, preferably 0.001% or more and 0.100% or less.
Zr: less than 0.200%
Zr is an effective element for spheroidizing the shape of sulfide and improving the adverse effect of sulfide on bendability, and thus may be contained as needed. However, excessive addition exceeding 0.200% causes an increase in inclusions and the like, causing surface and internal defects and the like. Therefore, when Zr is added, the addition amount is set to 0.200% or less, preferably to 0.0005% or more and 0.200% or less.
Ca: less than 0.0050%, mg: less than 0.0050%, REM: less than 0.0050%
Ca. Mg and REM are effective elements for spheroidizing the shape of sulfide and improving the adverse effect of sulfide on hole expansibility, and therefore may be contained as needed. However, excessive addition of each exceeding 0.0050% causes an increase in inclusions and the like, and causes surface and internal defects and the like. Therefore, when Ca, mg, and REM are added, the addition amounts thereof are each set to 0.0050% or less, preferably to 0.0005% or more and 0.0050% or less.
The balance other than the above components is Fe and unavoidable impurities.
(2) Next, the steel structure will be described.
Area ratio of ferrite: 1% to 40% inclusive
In order to ensure sufficient ductility, it is necessary to set the area ratio of ferrite to 1% or more. In order to secure a TS of 980MPa or more, the area ratio of soft ferrite needs to be 40% or less. The ferrite herein refers to polygonal ferrite, granular ferrite, or acicular ferrite, and is relatively soft and ductile ferrite. Preferably 3% or more and 30% or less.
Area ratio of fresh martensite: 1% to 20%
In order to achieve TS of 980MPa or more, it is necessary to set the area ratio of fresh martensite to 1% or more. In order to ensure good hole expansibility, the area ratio of fresh martensite needs to be 20% or less. Preferably 3% or more and 18% or less.
The sum of the area ratio of the bainite and the tempered martensite is 35% or more and 90% or less
Bainite and tempered martensite are structures effective for improving hole expansibility. When the sum of the area ratios of bainite and tempered martensite is less than 35%, good hole expansibility cannot be obtained. Therefore, the sum of the area ratio of bainite and tempered martensite needs to be 35% or more. On the other hand, when the sum of the area ratios of bainite and tempered martensite is more than 90%, desired retained austenite responsible for ductility is not obtained, and thus good ductility is not obtained. Therefore, the sum of the area ratio of bainite and tempered martensite needs to be 90% or less. Preferably 45% to 85%.
The area ratios of ferrite, fresh martensite, tempered martensite, and bainite can be obtained as follows: after polishing a plate thickness cross section (L cross section) parallel to the rolling direction of the steel plate, the steel plate was corroded with a 3 vol% nitric acid/ethanol solution, 10 fields of view were observed at 2000 x magnification using SEM (scanning electron microscope) at a position 1/4 of the plate thickness (corresponding to a position 1/4 of the plate thickness in the depth direction from the surface of the steel plate), and the area ratio of each of the structures (ferrite, fresh martensite, tempered martensite, bainite) of 10 fields of view was calculated using Image-Pro of Media Cybernetics company, and the values were averaged. In the above-described microstructure image, ferrite is a gray microstructure (base microstructure), fresh martensite is a white microstructure, tempered martensite is a microstructure having a gray internal structure in white martensite, and bainite is a dark gray microstructure having a large number of linear grain boundaries.
Area ratio of retained austenite: more than 6 percent
In order to ensure sufficient ductility, the area ratio of the retained austenite needs to be 6% or more. Preferably 8% or more. More preferably 10% or more.
The area ratio of the retained austenite was measured by measuring the integrated intensity ratios of diffraction peaks of {200}, {220}, {311} plane and bcc iron {200}, {211}, and {220} plane of fcc iron using a cokα ray by an X-ray diffraction apparatus on the obtained surface by grinding the steel sheet to a surface 0.1mm from the plate thickness 1/4 position and then further grinding off 0.1mm by chemical grinding, and the obtained 9 integrated intensity ratios were averaged.
A value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite: 1.1 or more
A value of 1.1 or more, which is obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite, is an extremely important constituent element in the present invention. In order to ensure good ductility, the area ratio of the Mn-enriched stable retained austenite needs to be high. Preferably 1.2 or more.
The value obtained by dividing the average C content (mass%) in the retained austenite having an aspect ratio of 2.0 or more by the average C content (mass%) in the ferrite is 3.0 or more
A value obtained by dividing the average C content (mass%) in the retained austenite having an aspect ratio (major axis/minor axis) of 2.0 or more by the average C content (mass%) in the ferrite is an extremely important constituent element in the present invention. In order to ensure good bendability, the area ratio of the stable retained austenite enriched with C needs to be high. Preferably 5.0 or more. The upper limit of the aspect ratio of the retained austenite is not particularly limited, and may be preferably 20.0 or less.
Regarding the amounts of C and Mn in the retained austenite and ferrite, the distribution state of Mn in each phase of the rolling direction section at the 1/4 thick position of the plate can be quantified using FE-EPMA (Field Emission-Electron Probe Micro Analyzer: field Emission electron probe microanalyzer), and the result of analysis of the amounts of C and Mn of 30 retained austenite grains and 30 ferrite grains can be averaged to obtain the result.
To identify retained austenite from retained austenite and martensite, the same field of view was observed using SEM (Scanning Electron Microscope ) and EBSD (Electron Backscattered Diffraction, electron back scattering diffraction). Subsequently, the retained austenite in the SEM image was identified by the Phase Map (Phase Map) of EBSD. The aspect ratio of the retained austenite was calculated by drawing an ellipse circumscribed to the retained austenite grains using Photoshop elements and dividing the major axis length by the minor axis length.
The amount of C in the total retained austenite divided by T 0 The C content in the composition is 1.0 or more
The amount of C in the total retained austenite divided by T 0 A value of 1.0 or more in the amount of C in the composition is an extremely important constituent element in the present invention. T (T) 0 The composition is such that fcc and bcc have the same free energy at an arbitrary temperature, austenite is fcc, and ferrite and bainite are bcc. By making the C content ratio fcc in the whole retained austenite equal to T which is the free energy of bcc 0 The amount of C in the composition is high, and decomposition of the retained austenite during the plating treatment can be suppressed, and a desired retained austenite amount can be obtained. As a result, the reduction in ductility conventionally reduced by the plating treatment can be prevented, and good ductility can be ensured. Therefore, it is necessary to divide the amount of C in the whole retained austenite by T 0 The amount of C in the composition is 1.0 or more. Preferably 1.1 or more.
The amount of C in the total retained austenite is calculated by using an X-ray diffraction apparatus, a K.alpha.ray using Co, a displacement of a diffraction peak using the (220) plane, and the following expressions [1] and [2 ].
a＝1.7889×√2/sinθ…[1]
a＝3.578+0.033[C]+0.00095[Mn]…[2]
Here, in the expressions [1] and [2], a is a lattice constant (), and θ is a value (rad) obtained by dividing a diffraction peak angle of the (220) plane by 2. [2] Wherein [ M ] is the mass% of the element M in the total austenite. In the present invention, the mass% of the element M in the retained austenite is set to be the mass% of the steel as a whole.
In addition, regarding T 0 The amount of C in the composition can be calculated uniquely from the composition of the steel and its content using the comprehensive thermodynamic calculation software Thermo-Calc and using the database TCFE 7. Calculated T 0 The composition is calculated at the reheat temperature before immersion in the galvanising bath.
In addition, it is preferable that the average ratio of the length to diameter of the retained austenite is 3.0 or more, which is obtained by multiplying the average amount of Mn (mass%) in the retained austenite by the average amount of Mn (mass%) in the ferrite. In order to ensure good ductility, it is necessary to increase the area ratio of the stable retained austenite having a large aspect ratio and enriched in Mn. Preferably 4.0 or more. The upper limit value is preferably 20.0 or less.
Further, it is preferable that the area ratio of the bulk retained austenite divided by the area ratio of the whole retained austenite and the bulk fresh martensite is 0.5 or less. Since the bulk retained austenite has high stability due to the restraint from the surrounding crystal grains, martensitic transformation occurs in a high strain region during blanking, the hardness difference from the surrounding crystal grains increases, and hole expansibility may be deteriorated. Therefore, the area ratio of the bulk retained austenite is preferably 0.5 or less divided by the area ratio of the whole retained austenite and the bulk fresh martensite. Preferably 0.4 or less. The bulk retained austenite refers to austenite having an aspect ratio of less than 2.0. The average crystal grain size of the bulk retained austenite is not limited, but for example, an average crystal grain size of 3 μm or less can be considered. The average crystal grain size can be obtained by a conventionally known method, for example, by performing image analysis on a tissue image of bulk retained austenite captured by a Scanning Electron Microscope (SEM).
In the steel structure of the present invention, in addition to ferrite, fresh martensite, bainite, tempered martensite, and retained austenite, even if carbide such as pearlite or cementite is contained in the range of 10% or less in terms of area ratio, the effect of the present invention is not impaired.
The high-strength steel sheet may further have a zinc plating layer. The zinc plating layer may be an alloyed zinc plating layer subjected to an alloying treatment.
(3) Next, production conditions will be described.
Heating temperature of billet
Although not particularly limited, the heating temperature of the billet is preferably set to 1100 ℃ or higher and 1300 ℃ or lower. The precipitates present in the heating stage of the steel slab are present as coarse precipitates in the finally obtained steel sheet and do not contribute to the strength, and therefore, it is preferable to redissolve Ti and Nb-based precipitates deposited during casting. Therefore, the heating temperature of the billet is preferably set to 1100 ℃ or higher. In addition, from the viewpoint of removing defects such as bubbles and segregation in the surface layer of the steel slab, reducing cracks and irregularities in the surface of the steel sheet, and realizing a smoother surface of the steel sheet, the heating temperature of the steel slab is preferably set to 1100 ℃ or higher. On the other hand, when the heating temperature of the billet exceeds 1300 ℃, the oxide scale loss increases with an increase in the oxidation amount, and therefore, the heating temperature of the billet is preferably set to 1300 ℃ or lower. More preferably, the temperature is 1150 ℃ or higher and 1250 ℃ or lower.
The billet is preferably manufactured by a continuous casting method in order to prevent macrosegregation, but may be manufactured by an ingot casting method, a thin slab casting method, or the like. In addition to the conventional method of once cooling to room temperature and then reheating after the billet is manufactured, there may be used, without any problem, an energy-saving process such as direct rolling in which a billet is charged into a heating furnace in a state of a hot plate without cooling to room temperature, or immediately rolled after slightly maintaining the temperature. In addition, the slab is formed into a thin slab by rough rolling under normal conditions, but in the case where the heating temperature is low, it is preferable to heat the thin slab using a rod heater or the like before finish rolling, from the viewpoint of preventing a failure at the time of hot rolling.
Finish rolling outlet side temperature of hot rolling: 750 ℃ to 1000 DEG C
The heated billet is hot rolled by rough rolling and finish rolling to produce a hot rolled steel sheet. At this time, when the finish rolling temperature exceeds 1000 ℃, the amount of oxide (scale) formed increases rapidly, and the interface between the steel base and the oxide becomes rough, and the surface quality after pickling and cold rolling tends to deteriorate. In addition, if there is partially a residue of hot rolled scale or the like after pickling, it adversely affects ductility and hole expansibility. In addition, the crystal grain size may become excessively coarse, and the surface of the pressed product may be roughened during processing. On the other hand, when the finish rolling temperature is lower than 750 ℃, the rolling load increases, the reduction ratio of austenite in an unrecrystallized state increases, abnormal texture develops, in-plane anisotropy in the final product becomes remarkable, uniformity of the material (material stability) is impaired, and the ductility itself decreases. Therefore, it is necessary to set the finish rolling outlet side temperature of the hot rolling to 750 ℃ or higher and 1000 ℃ or lower. Preferably, the temperature is set to 800 to 950 ℃.
Coiling temperature after hot rolling: 300 ℃ to 750 DEG C
When the coiling temperature after hot rolling exceeds 750 ℃, the grain size of ferrite of the hot rolled sheet structure becomes large, and it is difficult to secure the desired strength of the final annealed sheet. On the other hand, when the coiling temperature after hot rolling is lower than 300 ℃, the strength of the hot rolled sheet increases, the rolling load during cold rolling increases, or defects in the sheet shape occur, and thus productivity decreases. Therefore, the coiling temperature after hot rolling needs to be set to 300 ℃ to 750 ℃. Preferably, the temperature is set to 400 ℃ to 650 ℃.
In the hot rolling, rough rolled plates may be joined to each other to continuously finish rolling. The rough rolled sheet may be temporarily wound. In order to reduce the rolling load during hot rolling, a part or the whole of the finish rolling may be lubrication-rolled. From the viewpoints of homogenization of the shape of the steel sheet and homogenization of the material quality, it is also effective to perform lubrication rolling. The friction coefficient at the time of lubrication rolling is preferably set to 0.10 or more and 0.25 or less.
The hot-rolled steel sheet thus produced is optionally pickled. Since the acid washing can remove oxides on the surface of the steel sheet, it is preferable to perform the acid washing in order to ensure good chemical conversion treatability and plating quality of the high-strength steel sheet of the final product. In the case of pickling, the pickling may be performed once or a plurality of times.
Cold rolling
After coiling, acid washing is performed as needed, and then cold rolling is performed. The cold rolling reduction is not particularly limited, but is preferably 5% to 60%.
At Ac 1 Maintained at a temperature below the transformation point for more than 1800 seconds
At Ac 1 The steel sheet which is maintained in a temperature range of 1800 seconds or less below the transformation point can be softened for use in the subsequent cold rolling, and thus can be performed as needed. At a temperature exceeding Ac 1 When the transformation point is maintained in a temperature range, mn is enriched in austenite, and after cooling, hard martensite and retained austenite are formed, and softening of the steel sheet may not be performed. In addition, when the steel sheet is held for 1800 seconds or less, the strain after hot rolling cannot be removed, and the steel sheet may not be softened.
The heat treatment method may be any annealing method, such as continuous annealing or batch annealing. The cooling method and cooling rate are not particularly limited, and any one of cooling by furnace cooling in batch annealing, air cooling, gas jet cooling in continuous annealing, spray cooling, water cooling, and the like may be used. In addition, in the case of carrying out the acid washing treatment, a conventional method may be employed.
At Ac 3 The temperature of the transformation point is maintained for 20s to 1800s (corresponding to the first annealing treatment of the cold-rolled sheet of the embodiment) in the temperature range of-50 DEG or more
At below Ac 3 The phase transition point is carried out within the temperature range of-50 DEG CIn the case of holding, mn is enriched in austenite, and martensitic transformation does not occur during cooling, so that a core of retained austenite having a large aspect ratio cannot be obtained. As a result, in the subsequent annealing step (corresponding to the second annealing treatment of the cold rolled sheet of example), retained austenite is formed from grain boundaries, and retained austenite having a small aspect ratio increases, and a desired structure is not obtained, so hole expansibility is reduced.
If the content is less than 20s, sufficient recrystallization is not performed, and a desired structure is not obtained, so that hole expansibility is lowered. In addition, the Mn surface enrichment for ensuring the quality of the plating layer thereafter is not sufficiently performed.
On the other hand, if the temperature is maintained over 1800 seconds, the Mn surface enrichment becomes excessive, the coating quality deteriorates, and not only does austenite grains during annealing coarsen, thereby the nuclei of the retained austenite formed during the subsequent cooling also coarsen, and C cannot be sufficiently enriched to T 0 Above the composition, ductility decreases after plating.
Cooling to a cooling stop temperature below the martensitic transformation start temperature
When the cooling stop temperature exceeds the martensite transformation start temperature, if the amount of martensite to be transformed is small, the non-transformed austenite is entirely transformed into martensite during the final cooling, and the core of the retained austenite having a large aspect ratio cannot be obtained. As a result, in the subsequent annealing step (corresponding to the second annealing treatment of the cold rolled sheet of example), retained austenite is formed from grain boundaries, and retained austenite having a small aspect ratio increases, and a desired structure is not obtained, so that ductility and hole expansibility are reduced.
Preferably, the martensite phase transition initiation temperature is not lower than-250 ℃ and not higher than-50 ℃.
Reheating to a reheating temperature in the range of 120 ℃ to 450 ℃ and below, maintaining the reheating temperature for 2s to 1800s, and cooling to room temperature
In the case of a reheating temperature of less than 120 ℃, C is not enriched in the residual austenite formed in the subsequent annealing step, and a desired structure is not obtained, so ductility, bendability, and ductility after plating are reduced. When the reheating temperature exceeds 450 ℃, the retained austenite having a large aspect ratio is decomposed into nuclei, and the retained austenite having a small aspect ratio increases, so that a desired structure is not obtained, and the ductility decreases. In addition, in the case of keeping the length to diameter ratio to be smaller than 2s, the core of the retained austenite having a large length to diameter ratio cannot be obtained, and the desired structure cannot be obtained, so that the ductility, bendability, and ductility after plating are lowered. In addition, if the retained content exceeds 1800 seconds, the retained austenite having a large aspect ratio is decomposed into nuclei, and the retained austenite having a small aspect ratio is increased, and thus a desired structure is not obtained, and the ductility is lowered.
After the reheating is maintained for a predetermined period of time, the mixture is cooled to room temperature. The cooling method is not particularly limited, and may be a known method.
At Ac 1 The temperature range of the transformation point above-20 ℃ is kept for 20s to 600s (corresponding to the second annealing treatment of the cold-rolled sheet of the embodiment)
At Ac 1 The invention is extremely important in the present invention in that the phase transition point is maintained at a temperature of 20 ℃ or higher for 20s or more and 600s or less. At below Ac 1 When the transformation point is maintained in a temperature range of-20 ℃ and less than 20 seconds, carbide formed during the temperature rise is not completely dissolved, and it is difficult to secure a sufficient area ratio of martensite and retained austenite, and the strength is lowered. Preferably Ac 1 Above the phase transition point. More preferably Ac 1 Transformation point +20deg.C and Ac 3 The phase transition point is less than +50℃. Further, in the case of holding for more than 600 seconds, austenite coarsens during annealing, so that diffusion of Mn into austenite becomes insufficient, enrichment becomes impossible, and a sufficient area ratio of retained austenite for securing ductility cannot be obtained.
Cooling to a cooling stop temperature below the martensitic transformation start temperature
When the cooling stop temperature exceeds the martensite transformation start temperature, the amount of martensite to be transformed is small, and the amount of tempered martensite is small in the subsequent reheating, so that the desired amount of tempered martensite is not obtained. Preferably, the martensite phase transition initiation temperature is not lower than-250 ℃ and not higher than-30 ℃.
Reheating to a reheating temperature in the range of 120 ℃ to 480 ℃ and then holding at the reheating temperature for 2s to 600s, and cooling to room temperature
In the case of reheating below 120 ℃, the fresh martensite is not tempered, and the desired structure is not obtained. At reheating temperatures exceeding 480 ℃, bainitic transformation is delayed, and the desired structure is not obtained. In addition, if the temperature is kept less than 2s, the bainite transformation does not proceed sufficiently, and thus a desired structure cannot be obtained. On the other hand, if the holding time exceeds 600 seconds, carbide precipitates during bainite transformation, and the amount of C in the retained austenite decreases, so that a desired structure cannot be obtained.
After being kept at this temperature for a predetermined period of time, the mixture was cooled to room temperature. The cooling method is not particularly limited, and may be a known method.
Galvanization treatment
The obtained high-strength steel sheet is subjected to a zinc plating treatment as needed. In the case of hot dip galvanizing, the annealed steel sheet is immersed in a zinc plating bath at 440 to 500 ℃ and then subjected to hot dip galvanizing, and then the plating amount is adjusted by gas wiping or the like. It is preferable to use a zinc plating bath having an Al content of 0.08% or more and 0.30% or less for hot dip galvanizing.
When the hot dip zinc coating is alloyed, the hot dip zinc coating is alloyed at a temperature of 450 to 600 ℃ inclusive after the hot dip zinc coating. When the alloying treatment is performed at a temperature exceeding 600 ℃, the non-transformed austenite is transformed into pearlite, and the desired area ratio of the retained austenite cannot be ensured, and the ductility may be lowered. Therefore, when the alloying treatment of the zinc plating layer is performed, the alloying treatment of the zinc plating layer is preferably performed at a temperature range of 450 ℃ or more and 600 ℃ or less.
The conditions of the other production methods are not particularly limited, and the annealing is preferably performed by a continuous annealing apparatus from the viewpoint of productivity. Further, a series of treatments such as annealing, galvanizing, and alloying of the galvanized layer are preferably performed by a hot galvanizing line CGL (continuous galvanizing line ).
The "high-strength steel sheet" and the "high-strength hot dip galvanized steel sheet" may be subjected to skin pass rolling for the purpose of shape correction, adjustment of surface roughness, and the like. The reduction ratio of skin pass rolling is preferably in the range of 0.1% to 2.0%. When the content is less than 0.1%, the effect is small and control is difficult, and therefore, the lower limit of the preferable range is defined. In addition, if it exceeds 2.0%, productivity is significantly lowered, and therefore it is regarded as the upper limit of the good range. The skin pass rolling may be performed on-line or off-line. The skin pass rolling of the target rolling reduction may be performed at one time or may be performed in a plurality of times. Various coating treatments such as resin and grease coating may be performed.
Examples
Steel having the composition shown in table 1 and the balance of Fe and unavoidable impurities was melted in a converter, and a billet was produced by a continuous casting method. After reheating the obtained steel slab to 1250 ℃, a high-strength cold-rolled steel sheet (CR) was obtained under the conditions shown in tables 2 and 3, and a hot-dip galvanized steel sheet (GI) and a galvannealed steel sheet (GA) were obtained by further performing a galvanization treatment. The plate thicknesses of CR, GI, and GA are 1.0mm to 1.8 mm. As the hot dip zinc plating bath, a hot dip zinc plated steel sheet (GI) containing Al: a zinc bath of 0.19 mass% was used for the galvannealed steel sheet (GA) containing Al:0.14 mass% zinc bath, the bath temperature was set to 465 ℃. The coating adhesion amount was set to 45g/m per single surface 2 GA is adjusted so that the Fe concentration in the plating layer is 9 mass% or more and 12 mass% or less (both-sided plating). The steel structure of the cross section of the steel sheet obtained by the above method was observed, and the tensile properties, hole expansibility, bendability, and plating properties were examined, and the results are shown in tables 4 to 6.
Martensitic transformation initiation temperature and Ac 1 Transformation point and Ac 3 The phase transition point is determined using the following equation.
Martensitic transformation onset temperature (DEGC) =550-350× (% C) -40× (% Mn) -10× (% Cu) -17× (% Ni) -20× (% Cr) -10× (% Mo) -35× (% V) -5× (% W) +30× (% Al)
Ac 1 Transformation point (°c) =751-16× (% C) +11× (% Si) -28× (% Mn) -5.5× (% Cu) -16× (% Ni) +13× (% Cr) +3.4× (% Mo)
Ac 3 Transformation point (°c) =910-203 × (%c) +45× (%si) -30× (%mn) -20× (%cu) -15× (%ni) +11× (%cr) +32× (%mo) +104× (%v) +400× (%ti) +200× (%al)
Here, (%c), (%si), (%mn), (%ni), (%cu), (%cr), (%mo), (%v), (%ti), (%w), (%al) are the contents (mass%) of the respective elements, and when not contained, the values were set to zero.
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TABLE 4
An underline section: indicating that it is outside the scope of the present invention.
F: ferrite, M: fresh martensite, RA: retained austenite
TM: tempered martensite, B: bainite
TABLE 5
An underline section: indicating that it is outside the scope of the present invention.
F: ferrite, RA: retained austenite, P: pearlite, θ: carbide (cementite, etc.)
TABLE 6
An underline section: indicating that it is outside the scope of the present invention.
In the tensile test, a JIS5 test piece was used in which a sample was cut so that the tensile direction was perpendicular to the rolling direction of the steel sheet, TS (tensile strength) and EL (total elongation) were measured according to JIS Z2241 (2011), and the post-plating ductility (EL/EL') was measured in the case of a plated steel sheet. Here, EL 'shows the total elongation when passing through the sheet without immersing in a plating bath, and in the case of a cold-rolled steel sheet, el=el'. The following was determined to be good regarding the mechanical properties.
TS: under 980MPa or more and 1180MPa or less, EL is not less than 20%, and EL/EL' is not less than 0.7
TS:1180MPa or more, EL is 12% or more, and EL/EL' is0.7% or more
Hole expansibility was carried out in accordance with JIS Z2256 (2010). After each of the obtained steel sheets was cut into 100mm×100mm, holes having a diameter of 10mm were punched out with a clearance of 12% ± 1%, and then, a punch having an inner diameter of 75mm was pressed into the holes with a beading force of 9 tons, the hole diameter at the time of occurrence of cracks was measured by pressing a punch having a 60 ° cone into the holes, and the limiting hole expansion ratio λ (%) was obtained according to the following formula, and hole expansibility was evaluated based on the value of the limiting hole expansion ratio.
Limiting hole expansion ratio λ (%) = { (D) f -D 0 )/D 0 }×100
Wherein D is f Pore diameter (mm) at the time of crack generation, D 0 Is the initial pore size (mm). In the present invention, the following is determined to be good for each TS range.
TS:980MPa or more and 1180MPa or less, lambda is not less than 15%
TS:1180MPa or more, lambda is 25% or more
For the bending test, a bending test piece having a width of 30mm and a length of 100mm was cut from each annealed steel sheet with the rolling direction as a bending axis (Bending direction), and the measurement was performed by the V-block method according to JIS Z2248 (1996). An n=3 test was performed at a pressing speed of 100 mm/sec at each bending radius, and the presence or absence of cracking was determined on the outside of the bending portion by a solid microscope, and the minimum bending radius at which cracking did not occur was defined as the limiting bending radius R (mm). In the present invention, it is determined that the bending property of the steel sheet is good when the limit bending R/t at 90V bending is not more than 2.5 (t: the thickness of the steel sheet).
The plating properties were evaluated by appearance. The case where the surface quality was properly ensured without appearance defects such as non-plating, uneven alloying, and other defects that impair the surface quality was judged as "good", the case where the surface quality was excellent without even uneven color, in particular, was judged as "good", the case where some slight defects were observed was judged as "delta", and the case where a large number of surface defects were observed was judged as "x". The excellent cases and the excellent cases are determined as the scope of the present invention.
The high-strength steel sheet of the present invention has a TS of 980MPa or more and is excellent in formability. On the other hand, in the comparative example, at least one of TS, EL, ductility after plating, λ, bendability, and plating property was poor.
Industrial applicability
According to the present invention, a high-strength steel sheet having excellent formability with TS (tensile strength) of 980MPa or more can be obtained. By applying the high-strength steel sheet of the present invention to, for example, an automobile structural member, improvement in fuel efficiency due to weight reduction of the automobile body can be achieved, and industrial utilization value is extremely high.
Claims (9)
1. A high strength steel sheet, comprising:
contains C in mass%: 0.030% or more and 0.250% or less, si:0.01% or more and 3.00% or less, mn:2.00% or more and 8.00% or less, P:0.100% or less, S: below 0.0200%, N: less than 0.0100%, al:0.001% or more and 2.000% or less, the balance being Fe and unavoidable impurities; and
A steel structure in which ferrite is 1% to 40% by area ratio, fresh martensite is 1% to 20% by area ratio, the sum of bainite and tempered martensite is 35% to 90% by area ratio, and retained austenite is 6% or more,
a value obtained by dividing the average Mn amount in mass% in the retained austenite by the average Mn amount in mass% in the ferrite is 1.1 or more, and a value obtained by dividing the average C amount in mass% in the retained austenite having an aspect ratio of 2.0 or more by the average C amount in mass% in the ferrite is 3.0 or more,
the amount of C in the total retained austenite divided by T 0 The amount of C in the composition is 1.0 or more.
2. The high-strength steel sheet according to claim 1, wherein the composition of the components further contains, in mass%, a composition selected from the group consisting of Ti: less than 0.200%, nb: less than 0.200%, V: less than 0.500%, W:0.500% or less, B: less than 0.0050%, ni: less than 1.000%, cr: less than 1.000%, mo: less than 1.000%, cu: less than 1.000%, sn:0.200% or less, sb: less than 0.200%, ta: less than 0.100%, zr: less than 0.200%, ca: less than 0.0050%, mg: less than 0.0050%, REM:0.0050% or less of at least one element.
3. The high-strength steel sheet according to claim 1 or 2, wherein the area ratio of bulk retained austenite divided by the area ratio of all retained austenite and bulk fresh martensite is 0.5 or less.
4. The high-strength steel sheet according to any one of claims 1 to 3, further comprising a zinc plating layer on the surface.
5. The high-strength steel sheet according to claim 4, wherein the zinc plating layer is an alloyed zinc plating layer.
6. A method for producing a high-strength steel sheet according to any one of claims 1 to 3, wherein,
a steel billet having the composition according to claim 1 or 2 is heated, hot-rolled under the condition that the finish rolling outlet side temperature is 750 ℃ or more and 1000 ℃ or less, coiled and cold-rolled under 300 ℃ or more and 750 ℃ or less, and then Ac 3 Maintaining the temperature of the transformation point at-50deg.C for 20s or more and 1800s or less, cooling to cooling stop temperature of martensite transformation start temperature or less, reheating to reheating temperature of 120 deg.C or more and 450 deg.C or less, maintaining the reheating temperature for 2s or more and 1800s or less, cooling to room temperature, and cooling to Ac 1 The temperature range of the transformation point-20 ℃ is maintained at 20-600 s, cooled to a cooling stop temperature of not higher than the martensitic transformation start temperature, reheated to a reheating temperature in the range of not lower than 120 ℃ and not higher than 480 ℃, maintained at the reheating temperature for 2-600 s, and cooled to room temperature.
7. The method for producing a high-strength steel sheet according to claim 6, wherein a galvanization treatment is further performed.
8. The method for producing a high-strength steel sheet according to claim 7, wherein the alloying treatment is performed at 450 ℃ or higher and 600 ℃ or lower after the galvanization treatment.
9. The method for producing a high-strength steel sheet according to any one of claims 6 to 8, wherein Ac is applied after the coiling and before the cold rolling 1 The temperature range below the phase transition point remains above 1800s.
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| JP2588420B2 (en) | 1988-04-11 | 1997-03-05 | 日新製鋼株式会社 | Method for producing ultra-high strength steel with good ductility |
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| JP6597811B2 (en) * | 2017-04-05 | 2019-10-30 | Ｊｆｅスチール株式会社 | High-strength cold-rolled steel sheet and manufacturing method thereof |
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