CN117651786A - Steel sheet, member, and method for producing same - Google Patents
Steel sheet, member, and method for producing same Download PDFInfo
- Publication number
- CN117651786A CN117651786A CN202280050129.6A CN202280050129A CN117651786A CN 117651786 A CN117651786 A CN 117651786A CN 202280050129 A CN202280050129 A CN 202280050129A CN 117651786 A CN117651786 A CN 117651786A
- Authority
- CN
- China
- Prior art keywords
- less
- steel sheet
- temperature
- content
- delayed fracture
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Pending
Links
- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 164
- 239000010959 steel Substances 0.000 title claims abstract description 164
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 18
- 239000002244 precipitate Substances 0.000 claims abstract description 49
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 44
- 229910052758 niobium Inorganic materials 0.000 claims abstract description 30
- 229910052757 nitrogen Inorganic materials 0.000 claims abstract description 11
- 239000012535 impurity Substances 0.000 claims abstract description 9
- 229910052698 phosphorus Inorganic materials 0.000 claims abstract description 6
- 229910052717 sulfur Inorganic materials 0.000 claims abstract description 6
- 238000001816 cooling Methods 0.000 claims description 72
- 238000000137 annealing Methods 0.000 claims description 58
- 238000010438 heat treatment Methods 0.000 claims description 45
- 238000005096 rolling process Methods 0.000 claims description 36
- 229910001566 austenite Inorganic materials 0.000 claims description 32
- 239000000203 mixture Substances 0.000 claims description 32
- 238000007747 plating Methods 0.000 claims description 23
- 238000000034 method Methods 0.000 claims description 22
- 229910052761 rare earth metal Inorganic materials 0.000 claims description 15
- 239000010960 cold rolled steel Substances 0.000 claims description 14
- 229910052799 carbon Inorganic materials 0.000 claims description 11
- 238000005097 cold rolling Methods 0.000 claims description 8
- 238000005304 joining Methods 0.000 claims description 8
- 230000009467 reduction Effects 0.000 claims description 7
- 229910052684 Cerium Inorganic materials 0.000 claims description 4
- 229910045601 alloy Inorganic materials 0.000 claims description 3
- 239000000956 alloy Substances 0.000 claims description 3
- 229910052720 vanadium Inorganic materials 0.000 claims description 3
- 229910052804 chromium Inorganic materials 0.000 claims description 2
- 238000007598 dipping method Methods 0.000 claims description 2
- 229910052802 copper Inorganic materials 0.000 claims 1
- 230000003111 delayed effect Effects 0.000 abstract description 102
- 239000002245 particle Substances 0.000 abstract description 18
- 230000000694 effects Effects 0.000 description 23
- 238000005452 bending Methods 0.000 description 21
- 229910001563 bainite Inorganic materials 0.000 description 20
- 238000012360 testing method Methods 0.000 description 17
- 229910000859 α-Fe Inorganic materials 0.000 description 17
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 15
- 230000015572 biosynthetic process Effects 0.000 description 14
- 230000000717 retained effect Effects 0.000 description 14
- 239000000243 solution Substances 0.000 description 12
- 239000010410 layer Substances 0.000 description 10
- 230000008569 process Effects 0.000 description 10
- 238000005098 hot rolling Methods 0.000 description 9
- 239000006104 solid solution Substances 0.000 description 9
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 8
- 230000007423 decrease Effects 0.000 description 8
- 229910052739 hydrogen Inorganic materials 0.000 description 8
- 239000001257 hydrogen Substances 0.000 description 8
- 239000002344 surface layer Substances 0.000 description 8
- 230000006872 improvement Effects 0.000 description 7
- 150000001247 metal acetylides Chemical class 0.000 description 7
- MUBZPKHOEPUJKR-UHFFFAOYSA-N Oxalic acid Chemical compound OC(=O)C(O)=O MUBZPKHOEPUJKR-UHFFFAOYSA-N 0.000 description 6
- 238000005204 segregation Methods 0.000 description 6
- 238000002791 soaking Methods 0.000 description 6
- 238000009864 tensile test Methods 0.000 description 6
- 238000005275 alloying Methods 0.000 description 5
- 238000005336 cracking Methods 0.000 description 5
- 238000000227 grinding Methods 0.000 description 5
- 238000010008 shearing Methods 0.000 description 5
- 238000005728 strengthening Methods 0.000 description 5
- 239000000126 substance Substances 0.000 description 5
- 239000013585 weight reducing agent Substances 0.000 description 5
- 229910001567 cementite Inorganic materials 0.000 description 4
- 230000007797 corrosion Effects 0.000 description 4
- 238000005260 corrosion Methods 0.000 description 4
- 239000013078 crystal Substances 0.000 description 4
- 238000002149 energy-dispersive X-ray emission spectroscopy Methods 0.000 description 4
- 238000011156 evaluation Methods 0.000 description 4
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 3
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 3
- 230000002411 adverse Effects 0.000 description 3
- 238000006243 chemical reaction Methods 0.000 description 3
- 238000005520 cutting process Methods 0.000 description 3
- 230000007547 defect Effects 0.000 description 3
- 238000009826 distribution Methods 0.000 description 3
- 238000009713 electroplating Methods 0.000 description 3
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 3
- 239000000463 material Substances 0.000 description 3
- 150000004767 nitrides Chemical class 0.000 description 3
- 230000003647 oxidation Effects 0.000 description 3
- 238000007254 oxidation reaction Methods 0.000 description 3
- 238000007670 refining Methods 0.000 description 3
- 238000003303 reheating Methods 0.000 description 3
- 238000005496 tempering Methods 0.000 description 3
- 238000003466 welding Methods 0.000 description 3
- 238000004804 winding Methods 0.000 description 3
- 229910021578 Iron(III) chloride Inorganic materials 0.000 description 2
- 238000002441 X-ray diffraction Methods 0.000 description 2
- 230000001133 acceleration Effects 0.000 description 2
- SOIFLUNRINLCBN-UHFFFAOYSA-N ammonium thiocyanate Chemical compound [NH4+].[S-]C#N SOIFLUNRINLCBN-UHFFFAOYSA-N 0.000 description 2
- 238000004458 analytical method Methods 0.000 description 2
- 239000003153 chemical reaction reagent Substances 0.000 description 2
- 230000006866 deterioration Effects 0.000 description 2
- 238000000921 elemental analysis Methods 0.000 description 2
- 238000005516 engineering process Methods 0.000 description 2
- 238000005246 galvanizing Methods 0.000 description 2
- 238000010191 image analysis Methods 0.000 description 2
- RBTARNINKXHZNM-UHFFFAOYSA-K iron trichloride Chemical compound Cl[Fe](Cl)Cl RBTARNINKXHZNM-UHFFFAOYSA-K 0.000 description 2
- 229910052746 lanthanum Inorganic materials 0.000 description 2
- 239000007788 liquid Substances 0.000 description 2
- 238000002844 melting Methods 0.000 description 2
- 230000008018 melting Effects 0.000 description 2
- 229910052751 metal Inorganic materials 0.000 description 2
- 239000002184 metal Substances 0.000 description 2
- 238000000465 moulding Methods 0.000 description 2
- 230000003287 optical effect Effects 0.000 description 2
- 235000006408 oxalic acid Nutrition 0.000 description 2
- OXNIZHLAWKMVMX-UHFFFAOYSA-N picric acid Chemical compound OC1=C([N+]([O-])=O)C=C([N+]([O-])=O)C=C1[N+]([O-])=O OXNIZHLAWKMVMX-UHFFFAOYSA-N 0.000 description 2
- 238000005498 polishing Methods 0.000 description 2
- 238000001556 precipitation Methods 0.000 description 2
- 238000003825 pressing Methods 0.000 description 2
- 239000002994 raw material Substances 0.000 description 2
- 229920006395 saturated elastomer Polymers 0.000 description 2
- 230000001629 suppression Effects 0.000 description 2
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 2
- 229910052726 zirconium Inorganic materials 0.000 description 2
- 229910018072 Al 2 O 3 Inorganic materials 0.000 description 1
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 1
- 229910001335 Galvanized steel Inorganic materials 0.000 description 1
- 229910052765 Lutetium Inorganic materials 0.000 description 1
- GRYLNZFGIOXLOG-UHFFFAOYSA-N Nitric acid Chemical compound O[N+]([O-])=O GRYLNZFGIOXLOG-UHFFFAOYSA-N 0.000 description 1
- 239000007864 aqueous solution Substances 0.000 description 1
- 230000008901 benefit Effects 0.000 description 1
- 239000011247 coating layer Substances 0.000 description 1
- 239000002131 composite material Substances 0.000 description 1
- 239000000470 constituent Substances 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 230000000593 degrading effect Effects 0.000 description 1
- 238000003618 dip coating Methods 0.000 description 1
- NWIPPCMREQXKRU-UHFFFAOYSA-N ethanol nitrate Chemical compound CCO.[O-][N+]([O-])=O NWIPPCMREQXKRU-UHFFFAOYSA-N 0.000 description 1
- 239000008397 galvanized steel Substances 0.000 description 1
- 238000007429 general method Methods 0.000 description 1
- 238000007654 immersion Methods 0.000 description 1
- 239000004615 ingredient Substances 0.000 description 1
- 230000009545 invasion Effects 0.000 description 1
- 229910052747 lanthanoid Inorganic materials 0.000 description 1
- 150000002602 lanthanoids Chemical class 0.000 description 1
- FZLIPJUXYLNCLC-UHFFFAOYSA-N lanthanum atom Chemical compound [La] FZLIPJUXYLNCLC-UHFFFAOYSA-N 0.000 description 1
- OHSVLFRHMCKCQY-UHFFFAOYSA-N lutetium atom Chemical compound [Lu] OHSVLFRHMCKCQY-UHFFFAOYSA-N 0.000 description 1
- 230000014759 maintenance of location Effects 0.000 description 1
- 238000000691 measurement method Methods 0.000 description 1
- 238000010297 mechanical methods and process Methods 0.000 description 1
- 230000005226 mechanical processes and functions Effects 0.000 description 1
- 230000007246 mechanism Effects 0.000 description 1
- 150000002739 metals Chemical class 0.000 description 1
- 229910017604 nitric acid Inorganic materials 0.000 description 1
- 229910052760 oxygen Inorganic materials 0.000 description 1
- 230000002093 peripheral effect Effects 0.000 description 1
- 238000012545 processing Methods 0.000 description 1
- 239000000047 product Substances 0.000 description 1
- 230000002035 prolonged effect Effects 0.000 description 1
- 230000001737 promoting effect Effects 0.000 description 1
- 238000011002 quantification Methods 0.000 description 1
- 238000010791 quenching Methods 0.000 description 1
- 230000000171 quenching effect Effects 0.000 description 1
- 238000001953 recrystallisation Methods 0.000 description 1
- 229910052706 scandium Inorganic materials 0.000 description 1
- SIXSYDAISGFNSX-UHFFFAOYSA-N scandium atom Chemical compound [Sc] SIXSYDAISGFNSX-UHFFFAOYSA-N 0.000 description 1
- 230000000087 stabilizing effect Effects 0.000 description 1
- 150000004763 sulfides Chemical class 0.000 description 1
- 230000003746 surface roughness Effects 0.000 description 1
- 229910052727 yttrium Inorganic materials 0.000 description 1
- VWQVUPCCIRVNHF-UHFFFAOYSA-N yttrium atom Chemical compound [Y] VWQVUPCCIRVNHF-UHFFFAOYSA-N 0.000 description 1
Classifications
-
- 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
-
- 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
- C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
-
- 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
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
-
- 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
- C21D8/021—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular fabrication or treatment of ingot or slab
-
- 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
- C21D8/0236—Cold rolling
-
- 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
- C21D8/0242—Flattening; Dressing; Flexing
-
- 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
-
- 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
-
- 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
- C21D8/0278—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular surface treatment
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
-
- 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
-
- 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
- C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
-
- 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
- C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
-
- 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
- C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
-
- 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
-
- 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
-
- 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
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
-
- 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
-
- 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
- C23C2/02—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
-
- 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
-
- 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
- C23C2/26—After-treatment
- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
-
- 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
- C23C2/26—After-treatment
- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
- C23C2/29—Cooling or quenching
-
- 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
- C23C2/34—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
-
- 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
- C21D2211/004—Dispersions; Precipitations
-
- 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
Landscapes
- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Chemical Kinetics & Catalysis (AREA)
- Oil, Petroleum & Natural Gas (AREA)
- Heat Treatment Of Sheet Steel (AREA)
Abstract
The invention provides a steel sheet having high strength and excellent delayed fracture resistance, and a method for producing the same. The steel sheet comprises in mass percentThere is C:0.15 to 0.45 percent of Si: less than 1.5%, mn: greater than 1.7%, P: less than 0.03%, S: less than 0.0040%, sol.al:0.20% or less, N: less than 0.005%, B:0.0015 to 0.0100 percent, and one or more of Nb and Ti: the total of the components is 0.005-0.080%, the balance is Fe and unavoidable impurities, the area ratio of martensite relative to the whole structure is 95-100%, the primary gamma particle diameter is less than 11.2 mu m, the number density A of the precipitate with the equivalent circular diameter of more than 500nm satisfies the formula: a (number/mm) 2 )≤8.5×10 5 ×[B]。
Description
Technical Field
The present invention relates to a steel sheet such as a high-strength steel sheet for cold press forming used for cold press forming in automobiles and the like, a member using the steel sheet, and a method for producing the same.
Background
In recent years, for the purpose of weight reduction and collision safety of automobiles, steel sheets having a tensile strength TS of 1310MPa or more have been used for automobile frame members. In addition, steel sheets having tensile strength TS of 1470MPa or more are being used for bumpers, impact beam members, and the like.
When a high-strength steel sheet having a tensile strength TS of 1470MPa or more is formed into a member by cold press forming, there is a possibility that delayed fracture occurs due to an increase in residual stress in the member and deterioration in delayed fracture resistance of the steel sheet itself.
Here, delayed fracture refers to the following phenomenon: when the member is placed in a hydrogen-introduced environment in a state where high stress is applied to the member, hydrogen is introduced into the steel sheet constituting the member, and the interatomic bonding force is reduced, and local deformation is generated, whereby micro cracks are generated, and the micro cracks progress to cause breakage.
As a technique for improving such delayed fracture characteristics, for example, a technique for improving delayed fracture characteristics by reducing coarse precipitates serving as starting points of delayed fracture failure is disclosed in patent document 1, which has a steel sheet with improved delayed fracture characteristics in terms of mass% contains C:0.13% or more and 0.40% or less, si: less than 1.5%, mn:1.8% or more and 4% or less, P: less than 0.02%, S: less than 0.0010%, sol.al: less than 0.2%, N: less than 0.0060%, B:0.0003% or more and less than 0.0035%, O: less than 0.0020% and also to satisfy [%Ti ]+[％Nb]＞0.007、[％Ti]×[％Nb] 2 ＜7.5×10 -6 The mode of (1) contains Nb:0.002% or more and less than 0.035% and Ti: one or two of 0.002% to less than 0.040%, the balance being Fe and unavoidable impurities, the alloy comprises 4000 carbides containing Fe as a main component, wherein the area ratio of martensite and bainite to the whole structure is set to be more than 90% and 100% or less, the average grain size of prior austenite grains is set to be 6-15 mu m, inclusion groups satisfying specific conditions are reduced, the aspect ratio is 2.0 or less and the long axis is 0.30-2 mu m, and 4000 carbides/mm are formed from Fe as a main component 2 The following steel structure has a plate thickness of 0.5 to 2.6mm and a tensile strength of 1320MPa or more.
Patent document 2 discloses a high-strength cold-rolled steel sheet having excellent hydrogen embrittlement resistance and workability, which is characterized by comprising, in mass%, C:0.05 to 0.30 percent of Si: below 2.0% (including 0%), mn: greater than 0.1% and 2.8% or less, P:0.1% or less, S: less than 0.005%, N: less than 0.01%, al:0.01 to 0.50% by weight or less and satisfies [%C by 0.01% or more in total]-[％Nb]/92.9×12-[％Ti]/47.9×12-[％Zr]91.2X12 > 0.03, and has a structure comprising tempered martensite in an area ratio of 50% or more (including 100%) and ferrite in a balance, wherein precipitates having an equivalent circular diameter of 1 to 10nm are contained per 1 μm 2 Precipitates having 20 or more tempered martensite and an equivalent circle diameter of 20nm or more, i.e., precipitates containing one or two or more of Nb, ti and Zr, are present per 1 μm 2 The tempered martensite is 10 or less, and the average grain size of ferrite surrounded by large-angle grain boundaries having a crystal orientation difference of 15 DEG or more is 5 [ mu ] m or less.
Prior art literature
Patent literature
Patent document 1: japanese patent No. 6354921
Patent document 2: japanese patent No. 4712882
Disclosure of Invention
Problems to be solved by the invention
However, the conventional technology is not sufficient as a technology for securing tensile strength TS of 1470MPa or more and having excellent delayed fracture resistance.
The present invention has been made to solve the above problems, and an object of the present invention is to provide a steel sheet and a member having a tensile strength of 1470MPa or more (TS. Gtoreq.1470 MPa) and excellent delayed fracture resistance, and a method for producing the same.
The excellent delayed fracture resistance is determined to have excellent delayed fracture resistance by the following evaluation.
(1) First, a strip-shaped test piece having a rolling direction of 100mm and a rolling direction of 30mm was cut from a position 1/4 of the coil width from the end in the width direction of the steel sheet (coil) obtained.
(2) The cutting of the end face on the long side having a length of 100mm was set as shearing, and bending was performed so that the burrs were on the outer peripheral side of the bend in the sheared state (without performing the mechanical processing for removing burrs), and the test piece was fixed by a bolt while maintaining the shape of the test piece at the time of bending.
The clearance in the shearing process was set to 13%, the rake angle was set to 1 °, and the bending process was performed such that the bending radius of the tip was 10mm and the angle inside the bending apex was 90 degrees (V-bend).
The punch used was a punch having a U-shape (the portion of the tip R is a semicircle and the thickness of the punch body is 2R) and having the same tip radius as the tip bending radius R, and the die used was a die having a corner R of 30 mm. Then, the depth of the punch pressing the steel sheet was adjusted, and the steel sheet was formed so that the bending angle (angle inside the bending apex) of the tip was 90 degrees (V-shape).
The test piece was fastened by a hydraulic jack so that the distance between the flange ends of the straight piece portions at the time of bending was the same as that at the time of bending (so that the opening of the straight piece portion due to springback was eliminated), and the bolt was fastened in this state. The bolts were fixed through holes of an oval shape (short axis 10mm, long axis 15 mm) provided in advance 10mm inward from the short side edges of the strip-shaped test piece.
(3) The obtained test piece after bolt tightening was immersed in a solution of 0.1 mass% ammonium thiocyanate in water and McIlvaine buffer at 1:1 and adjusting the pH to 8.0, and performing a delayed fracture resistance evaluation test. At this time, the temperature of the solution was set to 20℃per 1cm of the test piece 3 The liquid amount of the surface area was set to 20ml.
(4) After 24 hours, it was confirmed whether or not there was a visually identifiable level (length of 1mm or more) of cracking, and it was determined that the delayed fracture resistance was excellent when no cracking was observed.
Means for solving the problems
The present inventors have conducted intensive studies to solve the above problems, and as a result, have found that the delayed fracture resistance can be significantly improved by satisfying all of the following conditions.
i) The area ratio of martensite is 95% or more.
ii) the prior austenite grains have an average grain size (prior gamma grain size) of less than 11.2 mu m.
iii) The number density A of precipitates having an equivalent circle diameter of 500nm or more satisfies the following condition.
A (number/mm) 2 )≤8.5×10 5 ×[B]
Here, [ B ] represents the content (mass%) of B.
The present invention has been further studied based on the above findings, and the gist thereof is as follows.
[1] A steel sheet comprising, in mass%, C:0.15% or more and 0.45% or less, si: less than 1.5%, mn: greater than 1.7%, P: less than 0.03%, S: less than 0.0040%, sol.al:0.20% or less, N: less than 0.005%, B:0.0015% or more and 0.0100% or less, and one or more of Nb and Ti: a total of 0.005% to 0.080%, and the balance of Fe and unavoidable impurities,
And has a structure in which the area of martensite relative to the whole structure is 95% or more and 100% or less,
the average grain size of the prior austenite grains is less than 11.2 mu m,
the number density A of precipitates having an equivalent circle diameter of 500nm or more satisfies the following formula (1).
A (number/mm) 2 )≤8.5×10 5 ×[B]… (1)
Here, [ B ] represents the content (mass%) of B.
[2] The steel sheet according to claim 1, further comprising, as the above-mentioned composition, a composition selected from the group consisting of, in mass%, cu:1.0% below and Ni:1.0% or less of one or two of the following.
[3] The steel sheet according to [1] or [2], wherein the steel sheet further comprises, as the above-mentioned composition, a composition selected from the group consisting of Cr: less than 1.0%, mo: less than 0.3%, V: less than 0.5%, zr:0.2% below and W:0.2% or less of one or two or more kinds of the above-mentioned components.
[4] The steel sheet according to any one of [1] to [3], wherein the steel sheet further comprises, as the above-mentioned composition, a composition selected from the group consisting of, in mass%, ca: less than 0.0030 percent, ce: less than 0.0030 percent, la: below 0.0030%, REM (excluding Ce, la): 0.0030% below and Mg:0.0030% or less of one or two or more kinds of the following.
[5] The steel sheet according to any one of [1] to [4], wherein the steel sheet further comprises, as the above-mentioned composition, a composition selected from the group consisting of, in mass%, sb:0.1% below and Sn:0.1% or less of one or two of the following.
[6] The steel sheet according to any one of [1] to [5], wherein a plating layer is provided on the surface of the steel sheet.
[7] A member obtained by using the steel sheet according to any one of [1] to [6 ].
[8] A method for manufacturing a steel sheet, wherein,
heating a steel billet having the composition of any one of [1] to [5] at an average heating rate of 10 ℃/min or less from 1000 ℃ to a heating holding temperature of 1250 ℃ or higher by a billet surface thermometer, holding at the heating holding temperature for 30 minutes or more, and then,
the residence time at 900-1000 ℃ is set to 20-150 seconds, the finish rolling is performed under the condition that the finish rolling temperature is set to 850 ℃ or higher,
cooling is performed at an average cooling rate of 40 ℃/sec or more in a range from the finish rolling temperature to 650 ℃,
then, coiling is performed at a coiling temperature of 650 ℃ or less, thereby producing a hot rolled steel sheet,
cold-rolling the hot-rolled steel sheet at a reduction of 40% or more to produce a cold-rolled steel sheet,
the following continuous anneals were performed:
the annealing temperature is set to 800-950 ℃, the cold-rolled steel sheet is heated from 400 ℃ to the annealing temperature at an average heating rate of 1.0 ℃/sec or more,
Kept at the annealing temperature for 600 seconds or less,
cooling from the annealing temperature to 420 ℃ at a first average cooling rate of more than 2 ℃/sec,
cooling from 420 ℃ to a cooling stop temperature below 260 ℃ at a second average cooling rate of more than 10 ℃/sec,
then, the mixture is kept at a holding temperature of 150 to 260 ℃ for 20 to 1500 seconds.
[9] The method of producing a steel sheet according to item [8], wherein the cold-rolled steel sheet is immersed in a plating bath while being cooled at the first average cooling rate in the continuous annealing, and is then heated to 480 to 600 ℃ after the immersion in the plating bath, whereby an alloying treatment is performed.
[10] The method of producing a steel sheet according to [8], wherein the surface of the steel sheet is subjected to a plating treatment after the continuous annealing.
[11] A method for producing a member, comprising the step of forming a steel sheet according to any one of [1] to [6] or performing at least one of forming and joining.
Effects of the invention
According to the present invention, a steel sheet and a member having high strength and excellent delayed fracture resistance and a method for producing the same can be provided.
Detailed Description
Hereinafter, embodiments of the present invention will be described.
The steel sheet of the present invention comprises, in mass%, C:0.15% or more and 0.45% or less, si: less than 1.5%, mn: greater than 1.7%, P: less than 0.03%, S: less than 0.0040%, sol.al:0.20% or less, N: less than 0.005%, B:0.0015% or more and 0.0100% or less, and one or more of Nb and Ti: the total composition is 0.005% to 0.080%, the balance is Fe and unavoidable impurities, and the composition has a structure in which the area ratio of martensite to the whole structure is 95% to 100%, the average particle diameter (primary gamma particle diameter) of the prior austenite grains (hereinafter also referred to as primary gamma grains) is less than 11.2 [ mu ] m, and the number density A of the precipitate having an equivalent circular diameter of 500nm or more satisfies the following formula (1).
A (number/mm) 2 )≤8.5×10 5 ×[B]… (1)
Here, [ B ] represents the content (mass%) of B.
Composition of the ingredients
The reasons for limiting the range of the composition of the steel sheet of the present invention will be explained below. The term "mass%" refers to the% of the component content.
C:0.15% or more and 0.45% or less
C is contained in order to increase the hardenability and to obtain a martensitic steel structure and to increase the strength of martensite. In order to secure tensile strength of 1470MPa or more (hereinafter, also referred to as TS. Gtoreq.1470 MPa), the C content is set to 0.15% or more. From the viewpoint of weight reduction of the automobile skeleton member due to an increase in tensile strength, the C content is preferably 0.20% or more, more preferably 0.27% or more. On the other hand, excessive addition of C becomes a factor that deteriorates delayed fracture resistance due to the formation of iron carbide and segregation to grain boundaries. From these viewpoints, the C content is limited to a range of 0.45% or less. The C content is preferably 0.40% or less, more preferably 0.37% or less.
Si: less than 1.5%
Si is contained as a strengthening element by solid solution strengthening and is contained from the viewpoint of improving the delayed fracture resistance by suppressing the formation of film-like carbide at the time of tempering in a temperature range of 200 ℃ or more. In addition, si is contained from the viewpoint of reducing Mn segregation at the plate thickness center portion and suppressing the formation of MnS. In addition, si is contained in order to suppress decarburization and B removal due to oxidation of the surface layer during annealing in a Continuous Annealing Line (CAL). The lower limit of the Si content is not specified, but from the viewpoint of obtaining the above-described effects, si is preferably contained at 0.02% or more. The Si content is preferably 0.10% or more, more preferably 0.20% or more. On the other hand, when the Si content is too large, the rolling load in hot rolling and cold rolling is significantly increased, and the toughness is lowered. From these viewpoints, the Si content is set to 1.5% or less (including 0%). The Si content is preferably 1.2% or less, more preferably 1.0% or less.
Mn: more than 1.7%
The Mn content is more than 1.7% in order to increase the hardenability of the steel and obtain a desired strength by setting the area ratio of martensite to a predetermined range. Preferably, the content is 1.8% or more. The upper limit is not particularly set, and Mn is an element that promotes the formation and coarsening of MnS particularly in the central portion of the plate thickness, and Al 2 O 3 The composite precipitation of inclusion particles such as (Nb, ti) (C, N), tiN, tiS and the like promotes delayed fracture. Therefore, the Mn content is preferably 4.0% or less. More preferably 3.0% or less.
P: less than 0.03%
P is an element for strengthening steel, and when the content thereof is large, segregation occurs at grain boundaries, and the grain boundary strength is lowered, so that the delayed fracture resistance and the spot weldability are remarkably deteriorated. From the above point of view, the P content is set to 0.03% or less. The P content is preferably 0.02% or less, more preferably 0.01% or less. The lower limit of the P content is not specified, but is set to 0.002% as a lower limit which is industrially practical at present.
S: less than 0.0040%
S forms coarse MnS and becomes the starting point of delayed fracture, thereby degrading delayed fracture resistance. From the viewpoint of improving the delayed fracture resistance, the S content needs to be set to at least less than 0.0040%. The S content is preferably less than 0.0020%, more preferably 0.0010% or less, and still more preferably 0.0007% or less. The lower limit is not specified, but is set to 0.0002% as a lower limit that can be industrially implemented at present.
sol.al: less than 0.20%
Al is contained in order to sufficiently deoxidize and reduce inclusions in steel. The lower limit of sol.al is not particularly limited, but in order to stably perform deoxidation, the content of sol.al is preferably set to 0.005% or more. The sol.al content is more preferably 0.01% or more, and still more preferably 0.02% or more. On the other hand, when the sol.al content exceeds 0.20%, cementite generated during coiling is difficult to be dissolved in solution during annealing, and the delayed fracture resistance is deteriorated. Therefore, the sol.al content is set to 0.20% or less. The sol.al content is preferably 0.10% or less, more preferably 0.05% or less.
N: less than 0.005%
N forms precipitates such as TiN, (Nb, ti) (C, N) in steel, and NbC, tiC, (Nb, ti) C, which are effective for reducing the prior austenite grain size, are reduced by the formation of these precipitates. They prevent the adjustment of the steel structure required for the present invention, and adversely affect the delayed fracture resistance. In order to reduce such adverse effects, the N content is set to 0.005% or less. The N content is preferably 0.0040% or less. The lower limit is not specified, but is set to 0.0006% as a lower limit which can be industrially implemented at present.
B:0.0015% or more and 0.0100% or less
B is an element that improves hardenability of steel, and has an advantage that martensite having a predetermined area ratio is generated even with a small Mn content. In addition, B increases the bonding force of the grain boundary due to segregation at the grain boundary, and suppresses segregation of P that decreases the grain boundary strength, thereby improving the delayed fracture resistance. On the other hand, the addition of excessive B gives Fe 23 (C,B) 6 BN increases and becomes a starting point of delayed fracture, resulting in a decrease in delayed fracture resistance. Therefore, in order to obtain the improvement of the delay break resistance due to the addition of BThe effect of the cracking characteristic needs to be achieved by both the increase in grain boundary solid solution B and the suppression of B-based precipitates. In the steel having an original gamma grain diameter of 10 μm or less, the B content is set to 0.0015% or more in order to obtain a sufficient grain boundary solid solution B content. The B content is preferably 0.0025% or more, more preferably 0.0040% or more. On the other hand, when B is contained in an amount of more than 0.0100%, it is difficult to reduce B-based precipitates even if hot rolling conditions and annealing conditions are controlled. Therefore, the B content is set to 0.0100% or less. The B content is preferably 0.0090% or less, more preferably 0.0080% or less.
One or more of Nb and Ti: the total of the components is more than 0.005% and less than 0.080%
Nb and Ti contribute to higher strength by miniaturization of the internal structure of martensite, and improve delayed fracture resistance by miniaturization of primary γ particle diameter. From such a viewpoint, at least one of Nb and Ti is contained in a total amount of 0.005% or more. The total content of Nb and Ti is preferably 0.010% or more, more preferably 0.020% or more. On the other hand, if the total of one or more of Nb and Ti exceeds 0.080%, nb and Ti are not completely solid-dissolved during reheating of the billet, precipitates having equivalent circle diameters of 500nm or more, such as TiN, ti (C, N), nbN, nb (C, N), nb (Nb, ti) (C, N), etc., increase, and become the starting points of delayed fracture, so that the delayed fracture resistance is rather deteriorated. Therefore, the upper limit of the total content of Nb and Ti is 0.080%. The total content of Nb and Ti (ti+nb) is preferably 0.07% or less, more preferably 0.06% or less.
The composition of the steel sheet of the present invention contains the above-mentioned constituent elements as essential components, and the balance contains iron (Fe) and unavoidable impurities. The steel sheet of the present invention preferably has a composition containing the above-described basic components, and the balance of iron (Fe) and unavoidable impurities.
In the present invention, one or two or more groups selected from the following (a) to (D) may be contained as the component composition.
(A) In mass%, selected from Cu:1.0% below and Ni:1.0% or less of one or two of the following;
(B) In mass%, selected from Cr: less than 1.0%, mo: less than 0.3%, V: less than 0.5%, zr:0.2% below and W: one or more than two of below 0.2%;
(C) In mass%, selected from Ca: less than 0.0030 percent, ce: less than 0.0030 percent, la: below 0.0030%, REM (excluding Ce, la): 0.0030% below and Mg:0.0030% or less of one or two or more kinds of the following materials;
(D) In mass%, selected from Sb:0.1% below and Sn:0.1% or less of one or two of the following.
Cu: less than 1.0%
Cu improves corrosion resistance in the environment where the automobile is used. Further, by containing Cu, the corrosion product coats the surface of the steel sheet, and the invasion of hydrogen into the steel sheet is suppressed. In addition, cu is an element mixed when scrap is effectively used as a raw material, and by allowing Cu to be mixed, a regenerated material can be effectively used as a raw material, and manufacturing cost can be reduced. From the above viewpoints, cu is preferably contained in an amount of 0.01% or more, and further from the viewpoint of improving the delayed fracture resistance, cu is preferably contained in an amount of 0.05% or more. The Cu content is more preferably 0.10% or more. However, if the content is too large, it becomes a cause of surface defects, and therefore the Cu content is preferably set to 1.0% or less. According to the above, when Cu is contained, the Cu content is set to 1.0% or less. The Cu content is more preferably 0.50% or less, and still more preferably 0.30% or less.
Ni: less than 1.0%
Ni is also an element having an effect of improving corrosion resistance. In addition, ni has an effect of reducing surface defects that are easily generated when Cu is contained. Therefore, from the above point of view, ni is preferably contained at 0.01% or more. The Ni content is more preferably 0.05% or more, and still more preferably 0.10% or more. However, when the Ni content is too large, the scale formation in the heating furnace becomes uneven, which causes surface defects and causes a significant cost increase. Therefore, when Ni is contained, the Ni content is set to 1.0% or less. The Ni content is more preferably 0.50% or less, and still more preferably 0.30% or less.
Cr: less than 1.0%
Cr may be added to obtain an effect of improving hardenability of the steel. In order to obtain this effect, cr is preferably contained in an amount of 0.01% or more. The Cr content is more preferably 0.05% or more, and still more preferably 0.10% or more. However, when the Cr content exceeds 1.0%, the solution rate of cementite at the time of annealing is retarded, and undissolved cementite remains, whereby the delayed fracture resistance of the sheared edge face is deteriorated. In addition, pitting resistance is also deteriorated. In addition, chemical conversion treatability is also deteriorated. Therefore, when Cr is contained, the Cr content is set to 1.0% or less. Since the delayed fracture resistance, pitting corrosion resistance and chemical conversion treatability all tend to deteriorate when the Cr content exceeds 0.2%, the Cr content is more preferably set to 0.2% or less from the viewpoint of preventing these.
Mo: less than 0.3%
Mo may be added for the purpose of obtaining an effect of improving hardenability of steel, an effect of forming fine carbides containing Mo as hydrogen trapping sites, and an effect of improving delayed fracture resistance by refining martensite. When Nb and Ti are added in large amounts, coarse precipitates are formed, and the delayed fracture resistance is rather deteriorated, but the solid solution limit amount of Mo is larger than that of Nb and Ti. When Nb and Ti are added in combination, fine precipitates are formed by combining them with Mo, and the effect of refining the structure is obtained. Therefore, by adding a small amount of Nb and Ti and adding Mo in combination, the microstructure can be made finer and the fine carbide can be dispersed in a large amount without leaving coarse precipitates, and the delayed fracture resistance can be improved. In order to obtain this effect, mo is preferably contained in an amount of 0.01% or more. The Mo content is more preferably 0.03% or more, and still more preferably 0.05% or more. However, when Mo is contained at 0.3% or more, chemical conversion treatability deteriorates. Therefore, in the case of containing Mo, the Mo content is set to less than 0.3%. The Mo content is preferably 0.2% or less.
V: less than 0.5%
V may be added for the purpose of obtaining an effect of improving hardenability of steel, an effect of forming fine carbides containing V as hydrogen trapping sites, and an effect of improving delayed fracture resistance by refining martensite. In order to obtain this effect, the V content is preferably set to 0.003% or more. The V content is more preferably 0.03% or more, and still more preferably 0.05% or more. However, when the V content exceeds 0.5%, the castability is remarkably deteriorated. Therefore, when V is contained, the V content is set to 0.5% or less. The V content is more preferably 0.3% or less, and still more preferably 0.2% or less. The V content is more preferably 0.1% or less.
Zr: less than 0.2%
Zr contributes to a higher strength and improves the delayed fracture resistance by the miniaturization of primary γ crystal grains and the resultant miniaturization of the internal structure of martensite. In addition, the strength is increased and the delayed fracture resistance is improved by forming fine Zr-based carbide/carbonitride serving as a hydrogen trapping site. In addition, zr improves castability. From such a viewpoint, the Zr content is preferably set to 0.005% or more. The Zr content is more preferably 0.010% or more, still more preferably 0.015% or more. However, when Zr is added in a large amount, coarse precipitates of ZrN and ZrS system remaining due to non-solid solution increase during heating of the slab in the hot rolling step, and the delayed fracture resistance of the sheared edge face is deteriorated. Therefore, when Zr is contained, the Zr content is set to 0.2% or less. The Zr content is more preferably 0.1% or less, and still more preferably 0.04% or less.
W: less than 0.2%
W contributes to a high strength and an improvement in delayed fracture resistance by forming fine W-based carbide and carbonitride serving as hydrogen trapping sites. From such a viewpoint, W is preferably contained in an amount of 0.005% or more. The W content is more preferably 0.010% or more, and still more preferably 0.030% or more. However, when W is contained in a large amount, coarse precipitates remaining due to non-solid solution increase during heating of a steel slab in the hot rolling step, and the delayed fracture resistance of the sheared edge face deteriorates. Therefore, when W is contained, the W content is set to 0.2% or less. The W content is more preferably 0.1% or less.
Ca: less than 0.0030 percent
Ca fixes S in the form of CaS to improve the delayed fracture resistance. In order to obtain this effect, ca is preferably contained in an amount of 0.0002% or more. The Ca content is more preferably 0.0005% or more, and still more preferably 0.0010% or more. However, when Ca is added in a large amount, the surface quality and bendability are deteriorated, and therefore, the Ca content is preferably 0.0030% or less. According to the above, when Ca is contained, the Ca content is set to 0.0030% or less. The Ca content is more preferably 0.0025% or less, and still more preferably 0.0020% or less.
Ce: less than 0.0030 percent
Ce also fixes S to improve delayed fracture resistance. In order to obtain this effect, ce is preferably contained at 0.0002% or more. The Ce content is more preferably 0.0003% or more, and still more preferably 0.0005% or more. However, when Ce is added in a large amount, the surface quality and bendability are deteriorated, and therefore the Ce content is preferably 0.0030% or less. According to the above, when Ce is contained, the Ce content is set to 0.0030% or less. The Ce content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.
La: less than 0.0030 percent
La also fixes S to improve delayed fracture resistance. In order to obtain this effect, la is preferably contained in an amount of 0.0002% or more. The La content is more preferably 0.0005% or more, and still more preferably 0.0010% or more. However, when a large amount of La is added, the surface quality and bendability are deteriorated, and therefore the La content is preferably 0.0030% or less. According to the above, when La is contained, the La content is set to 0.0030% or less. The La content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.
REM: less than 0.0030 percent
REM also fixes S to improve delayed fracture resistance. In order to obtain this effect, REM is preferably contained in an amount of 0.0002% or more. The REM content is more preferably 0.0003% or more, and still more preferably 0.0005% or more. However, when REM is added in a large amount, the surface quality and bendability are deteriorated, and therefore the REM content is preferably 0.0030% or less. According to the above, when REM is contained, the REM content is set to 0.0030% or less. The REM content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.
In the present invention, REM refers to scandium (Sc) having an atomic number of 21, yttrium (Y) having an atomic number of 39, and elements other than Ce and La in lanthanoids of lanthanum (La) having an atomic number of 57 to lutetium (Lu) having an atomic number of 71. In the present invention, REM concentration means the total content of one or more elements selected from the REM.
Mg: less than 0.0030 percent
Mg fixes O in the form of MgO to improve delayed fracture resistance. In order to obtain this effect, mg is preferably contained at 0.0002% or more. The Mg content is more preferably 0.0005% or more, and still more preferably 0.0010% or more. However, when Mg is added in a large amount, the surface quality and bendability are deteriorated, so that the Mg content is preferably 0.0030% or less. According to the above, when Mg is contained, the Mg content is set to 0.0030% or less. The Mg content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.
Sb: less than 0.1%
Sb suppresses oxidation and nitridation of the surface layer, and suppresses the decrease in C, B caused by this. By suppressing the decrease in C, B, ferrite generation in the surface layer is suppressed, which contributes to enhancement of strength and improvement of delayed fracture resistance. From such a viewpoint, the Sb content is preferably set to 0.002% or more. The Sb content is more preferably 0.004% or more, still more preferably 0.006% or more. However, if the Sb content exceeds 0.1%, the castability is deteriorated, and Sb segregates in the original γ grain boundary, and the delayed fracture resistance of the sheared edge face is deteriorated. Therefore, the Sb content is preferably 0.1% or less. According to the above, when Sb is contained, the Sb content is set to 0.1% or less. The Sb content is more preferably 0.05% or less, and still more preferably 0.02% or less.
Sn: less than 0.1%
Sn suppresses oxidation and nitridation of the surface layer, and suppresses the decrease in the content of C, B in the surface layer. By suppressing the decrease in C, B, ferrite generation in the surface layer is suppressed, which contributes to enhancement of strength and improvement of delayed fracture resistance. From such a viewpoint, the Sn content is preferably set to 0.002% or more. The Sn content is preferably 0.003% or more. However, when the Sn content exceeds 0.1%, the castability is deteriorated, and Sn segregates in the original γ grain boundary, and the delayed fracture resistance of the sheared edge face is deteriorated. Therefore, when Sn is contained, the Sn content is set to 0.1% or less. The Sn content is more preferably 0.05% or less, and still more preferably 0.01% or less.
In the case where the above-mentioned optional element is contained below the preferable lower limit value, it is set that the above-mentioned optional element is contained as an unavoidable impurity.
Steel structure
The steel structure of the steel sheet of the present invention has the following structure.
The area ratio of martensite to the whole structure is 95% to 100%.
(constitution 2) the prior austenite grains have an average grain diameter of less than 11.2. Mu.m.
(constitution 3) the number density A of precipitates having an equivalent circle diameter of 500nm or more satisfies the following formula (1).
A (number/mm) 2 )≤8.5×10 5 ×[B]… (1)
Here, [ B ] represents the content (mass%) of B.
Hereinafter, each configuration will be described.
The area ratio of martensite to the whole structure is 95% to 100%.
In order to achieve both high strength of TS not less than 1470MPa and excellent delayed fracture resistance, the area ratio of martensite in the steel structure is set to 95% or more. More preferably 99% or more, still more preferably 100%. When a structure other than martensite is contained, the balance may be bainite, ferrite, or retained austenite (retained γ). These tissues are also very small amounts of carbides, sulfides, nitrides, and oxides. The remaining amount of the tissue is 5% or less, preferably 1% or less. The martensite also includes martensite in which tempering including self-tempering during continuous cooling does not occur and which is performed by staying at about 150 ℃ or more for a certain period of time. The area ratio of martensite may be 100% without the remaining amount.
(constitution 2) the prior austenite grains have an average grain diameter of less than 11.2. Mu.m.
In steel having a martensite area ratio of 95% or more in the steel structure, the delayed fracture surface is mostly a grain boundary surface, and the start point of delayed fracture and the crack progress path at the early stage of delayed fracture are considered to be on the prior austenite grain boundary. In order to suppress the inter-grain fracture, the prior austenite grains are effectively refined, and the delayed fracture resistance is remarkably improved by the refinement of the prior austenite grains. As a mechanism, it is considered that: as the prior austenite grains are refined, the area ratio of the prior austenite grain boundaries increases, and the concentration of impurity elements such as P, which are grain boundary embrittling elements, on the prior austenite grain boundaries decreases. In addition, the refinement of the prior austenite grains contributes to the improvement of the tensile strength. From the viewpoints of delayed fracture resistance and strength, the prior austenite grains have an average grain size (prior γ grain size) of less than 11.2 μm. The average particle diameter is preferably 10 μm or less, more preferably 7.0 μm or less, and even more preferably 5.0 μm or less.
(constitution 3) the number density A of precipitates having an equivalent circle diameter of 500nm or more satisfies the following formula.
A (number/mm) 2 )≤8.5×10 5 ×[B]… (1)
Here, [ B ] represents the content (mass%) of B.
In a steel having a high strength of TS not less than 1470MPa, it is effective to segregate B at grain boundaries to strengthen the grain boundaries in addition to the refinement of prior austenite grains in order to suppress intergranular fracture. However, an increase in the amount of B alone increases not only grain boundary segregation B but also Fe as a starting point of delayed fracture 23 (C,B) 6 The B-based precipitates, which are the main components, also increase, and therefore the delayed fracture resistance is rather lowered. The inventors have found that by controlling the hot rolling conditions to reduce the number density A of precipitates having an equivalent circle diameter of 500nm or more and satisfying the following conditions, both improvement of delayed fracture resistance due to grain boundary strengthening of B and suppression of fracture at the starting point of the precipitates can be achieved.
A (number/mm) 2 )≤8.5×10 5 ×[B]
Preferably A (number/mm) 2 )≤5.0×10 5 ×[B]More preferably A (individual/mm 2 )≤2.0×10 5 ×[B]。
The measurement method of each structure in the above steel structure will be described.
The area ratios of martensite, bainite, and ferrite were measured as follows: after the L-section of the steel sheet (a section parallel to the rolling direction and perpendicular to the surface of the steel sheet (hereinafter also referred to as a perpendicular section parallel to the rolling direction)) was polished, the steel sheet was etched with an aqueous solution of nitric acid and observed for 4 fields of view at a magnification of 2000 times from the surface of the steel sheet by SEM, and the photographed tissue photograph was subjected to image analysis to determine the area ratio of martensite, bainite, and ferrite. Here, martensite and bainite refer to structures that are gray or white in SEM. On the other hand, ferrite is a region of contrast that appears black in SEM. Note that, although a small amount of carbide, nitride, sulfide, or oxide is contained in the martensite or bainite, it is difficult to exclude them, and therefore the area ratio of the region containing them is defined as the area ratio.
Here, bainite has the following characteristics. That is, the aspect ratio is 2.5 or more and the plate-like form is a slightly black structure compared with martensite. The width of the plate is 0.3-1.7 μm. The distribution density of carbide with diameter of 10-200 nm in bainite is 0-3/μm 2 。
The retained austenite (retained γ) is measured as follows: the surface layer of the steel sheet was subjected to chemical polishing with oxalic acid at 200 μm, and the retained austenite (retained γ) was obtained by an X-ray diffraction intensity method with respect to the sheet surface. Calculated from the integrated intensities of (200) α, (211) α, (220) α, (200) γ, (220) γ, and (311) γ diffraction plane peaks measured by Mo-kα rays.
The average grain size (primary γ grain size) of the prior austenite grains was measured as follows: after grinding the L-section (vertical section parallel to the rolling direction) of the steel sheet, the steel sheet was corroded with a reagent (for example, a saturated aqueous picric acid solution or a solution obtained by adding ferric chloride thereto) that corrodes the original γ grain boundaries, 4 fields of view were observed with an optical microscope at a magnification of 500 times at a position 1/4 the thickness from the steel sheet surface, and 15 lines were drawn at intervals of 10 μm or more in the actual length in the sheet thickness direction and the rolling direction in the obtained photographs, and the number of intersections of the grain boundaries and the lines was counted. Further, the prior γ grain size (average grain size of prior austenite grains) can be measured by multiplying a value obtained by dividing the number of intersections by the bus length by 1.13.
The number density A of precipitates having an equivalent circle diameter of 500nm or more was determined as follows: after grinding the L-section (vertical section parallel to the rolling direction) of the steel sheet, the steel sheet was continuously shot with SEM for 2mm in the region from 1/5 position to 4/5 position of the sheet thickness, i.e., the region from 1/5 position relative to the sheet thickness to 4/5 position from the center of the sheet thickness on the surface of the steel sheet 2 The number of such precipitates were measured from the SEM photograph, and the number density A of the precipitates having an equivalent circle diameter of 500nm or more was obtained. The magnification at which photographing was performed was 2000 times. In addition, when the composition analysis of each inclusion particle was performed, each inclusion particle was amplified 10000 times, and the above precipitate was analyzed. Here, the precipitate having an equivalent circle diameter of 500nm or more is Fe 23 (C,B) 6 And a precipitate containing B, and the presence or absence of a B peak was examined by elemental analysis using an energy dispersive X-ray spectrometry (EDS) with an acceleration voltage of 3kV, and the presence of a B peak was evaluated as the presence of the precipitate.
If the reheating of the billet is insufficient, precipitates containing Nb and Ti also increase, and these precipitates adversely affect the delayed fracture characteristics.
The equivalent circle diameter is the diameter of a perfect circle having the area of each precipitate calculated from SEM photographs.
Tensile Strength (TS): 1470MPa or more
The deterioration of the delayed fracture resistance is significantly remarkable when the tensile strength of the steel sheet is 1470MPa or more. It is one of the features of the present invention that the tensile strength is 1470MPa or more and the delayed fracture resistance is excellent. Therefore, in the present invention, the tensile strength is required to be 1470MPa or more. From the viewpoint of weight reduction of the automobile skeleton member, 1700MPa or more is preferable. The tensile strength of the steel sheet of the present invention may be 2100MPa or less.
The tensile strength can be determined as follows: a JIS No. 5 tensile test piece was cut at a position 1/4 of the coil width so that the direction of the roll was the longitudinal direction, and the tensile strength was measured by a tensile test according to JIS Z2241.
The steel sheet of the present invention may be a steel sheet having a plating layer on the surface thereof. The plating layer may be a Zn plating layer or a plating layer of other metals. The plating layer may be a hot-dip coating layer or a plating layer.
Next, a method for manufacturing the steel sheet of the present invention will be described.
The method for producing a steel sheet according to the present invention is a method for producing a steel sheet, wherein a steel slab having the above-mentioned composition is heated from 1000 ℃ to 1250 ℃ or higher at an average heating rate of 10 ℃/min or lower at a billet surface thermometer, and is held at the heating holding temperature for 30 minutes or more, then a residence time at 900 to 1000 ℃ is set to 20 seconds or more and 150 seconds or less, hot finish rolling is performed under conditions where a finish rolling temperature is set to 850 ℃ or more, cooling is performed where an average cooling rate in a range from the finish rolling temperature to 650 ℃ is set to 40 ℃/sec or more, and then coiling is performed at a coiling temperature of 650 ℃ or lower, thereby producing a hot rolled steel sheet, the hot rolled steel sheet is cold rolled at a reduction rate of 40% or more, thereby producing a cold rolled steel sheet, and the following continuous annealing is performed: the annealing temperature is set to 800-950 ℃, the cold-rolled steel sheet is heated from 400 ℃ to the annealing temperature at an average heating rate of 1.0 ℃/sec or more, kept at the annealing temperature for 600 seconds or less, cooled from the annealing temperature to 420 ℃ at a first average cooling rate of 2 ℃/sec or more, cooled from 420 ℃ to a cooling stop temperature of 260 ℃ or less at a second average cooling rate of 10 ℃/sec or more, and then kept at the keeping temperature of 150-260 ℃ for 20-1500 seconds.
Hot rolling
In heating a billet before hot rolling, the average heating rate of the heating maintaining temperature from 1000 ℃ to 1250 ℃ is set to 10 ℃/min or less, thereby promoting solid solution of sulfide and reducing the size and number of inclusions. Since the melting temperatures of Nb and Ti are high, the solid solution promotion of Nb and Ti can be achieved by setting the heating holding temperature to 1250 ℃ or higher and the holding time to 30 minutes or longer by the billet surface thermometer, and the size and number of precipitates can be reduced. The heating and holding temperature is preferably 1300 ℃ or higher. More preferably at 1350 ℃.
The average heating rate is herein defined as "(temperature at which the billet is heated (heating maintaining temperature) (°c) -temperature at which the billet is heated (°c) (1000 ℃))/heating time (minutes) from the start of heating to the completion of heating)".
Then, the billet is allowed to stand at 900 to 1000 ℃ for 20 seconds to 150 seconds. The increase of the residence time in the temperature range of 900 to 1000 ℃ causes the formation and coarsening of a precipitate mainly composed of BN. The precipitates formed in these temperature ranges are less likely to be solid-dissolved by annealing and heating, and the amount of solid-dissolved B after annealing is reduced. Therefore, when the residence time exceeds 150 seconds, the amount of solid solution B effective for suppressing delayed fracture cannot be obtained. Therefore, the residence time is 150 seconds or less, preferably 120 seconds or less, and more preferably 100 seconds or less. On the other hand, when the residence time is less than 20 seconds, the tissue may become uneven. Therefore, the residence time is 20 seconds or more, preferably 30 seconds or more, and more preferably 40 seconds or more.
In the hot finish rolling, the finish rolling temperature (FT) is set to 850 ℃ or higher in order to suppress precipitation of Nb, ti, B, etc. Preferably, the finish rolling temperature is 930 ℃ or lower.
In the cooling after the finish hot rolling, the average cooling rate in the range from the finish rolling temperature to 650 ℃ is set to 40 ℃/sec or more. When the average cooling rate is less than 40 ℃/sec, carbonitrides having an equivalent circle diameter of 1.0 μm or more increase due to coarsening of Nb carbonitrides and Ti carbonitrides, and the desired delayed fracture resistance cannot be obtained. The average cooling rate is preferably 250 ℃/sec or less, more preferably 200 ℃/sec or less.
The average cooling rate in the hot rolling step means "(temperature at the start of cooling (finish rolling temperature) (°c) -temperature at the end of cooling (°c) (650 ℃))/cooling time (seconds) from the start of cooling to the end of cooling)".
After the cooling to 650 ℃, the material was cooled and wound as needed. At this time, when the winding temperature exceeds 650 ℃, only the Nb and Ti-based precipitates precipitated in the fine austenite region coarsen, and therefore, the coarse precipitates increase and the delayed fracture characteristics decrease. Therefore, the winding temperature is set to 650 ℃ or lower. The winding temperature is preferably 500 ℃ or higher.
Cold rolling
In cold rolling, if the rolling reduction (cold rolling reduction) is set to 40% or more, the recrystallization behavior and texture orientation in the subsequent continuous annealing can be stabilized. If the content is less than 40%, part of austenite grains during annealing may become coarse, and the strength may be lowered. The cold rolling rate is preferably 80% or less.
Continuous annealing
The cold rolled steel sheet is annealed in a Continuous Annealing Line (CAL), tempered, and temper rolled as necessary.
Fe 23 (C,B) 6 Ferrite regions are formed and coarsened during annealing and heating, and therefore, in order to make Fe 23 (C,B) 6 It is extremely important to reduce the effect of grain boundary strengthening by B and to sufficiently increase the average heating rate at 400 ℃. In addition, from the viewpoint of reducing the primary γ particle diameter to less than 11.2 μm, it is also necessary to increase the heating rate. From the above, the average heating rate at 400℃or higher was 1.0℃or higher. The average heating rate at 400℃or higher is preferably 1.5℃or higher, more preferably 3.0℃or higher.
The average heating rate is preferably 10 ℃/sec or less.
The average heating rate herein means "annealing temperature (c) to 400 (c) described later"/heating time (minutes) from 400 to annealing temperature ".
To sufficiently reduce Fe remaining after annealing due to undissolved solution 23 (C,B) 6 And the like, and annealing is performed at a high temperature for a long time. Specifically, the annealing temperature needs to be set to 800 ℃ or higher.
On the other hand, in the case of annealing at a temperature exceeding 950 ℃, the primary γ particle diameter becomes coarse, and the target structure is not obtained, so the annealing temperature is set to 950 ℃ or lower. In addition, in the case of annealing at a temperature exceeding 900 ℃, BN is precipitated at grain boundaries, and the delayed fracture resistance is sometimes deteriorated, so that 900 ℃ or lower is more preferable. Since the raw γ grain size is too large even after the soaking time (holding time) at the annealing temperature is prolonged, the soaking time is set to 600 seconds or less. Preferably, the soaking time is 10 seconds or longer.
Then, in order to reduce ferrite and retained austenite and to make the area ratio of martensite to 95% or more, it is necessary to cool the steel sheet from the annealing temperature to 420 ℃ at a first average cooling rate of 2 ℃/sec or more. When the first average cooling rate is less than 2 ℃/sec, ferrite is largely generated, and carbon is enriched in γ, and martensite is hardened, so that the delayed fracture resistance is deteriorated. The upper limit of the first average cooling rate is not particularly limited, but is preferably 100 ℃/sec.
In the case of hot dip galvanizing a steel sheet, in the process of cooling from the annealing temperature to 420 ℃, more specifically, in the case of cooling at the above-mentioned first average cooling rate in continuous annealing, it is preferable to dip the cold-rolled steel sheet in a plating bath to perform a plating treatment, and if necessary, after dipping in the plating bath, it may be heated to 480 to 600 ℃ to perform an alloying treatment.
Next, in the present invention, in order to suppress the formation of bainitic ferrite and lower bainite and to make the area ratio of martensite be 95% or more, it is necessary to cool the bainitic ferrite and lower bainite from 420 ℃ to 260 ℃ or less at a second average cooling rate of 10 ℃/sec or more. In a structure in which bainite is largely formed, strength is lowered, and retained austenite is increased, so that delayed fracture resistance is deteriorated. Therefore, the second average cooling rate from 420 ℃ to 260 ℃ or lower cooling stop temperature is set to 10 ℃/sec or higher. The second average cooling rate is preferably 20 ℃/sec or more, more preferably 70 ℃/sec or more. The upper limit of the second average cooling rate is not particularly limited, but is preferably 2000 ℃/sec.
Herein, the first average cooling rate means "(annealing temperature (c) -420 (c))/(cooling time (seconds) from annealing temperature to 420 ℃).
In addition, the second average cooling rate means "(420 (DEG C) -cooling stop temperature (DEG C))/(cooling time (seconds) from 420 ℃ to cooling stop temperature).
When the cooling stop temperature exceeds 260 ℃, there is a problem that the formation of upper bainite/lower bainite, retained austenite, and fresh martensite increases. Therefore, the cooling stop temperature is set to 260 ℃ or lower.
The carbide distributed in the martensite is a carbide generated in the low-temperature range after quenching. In order to ensure excellent delayed fracture resistance and tensile strength of 1470MPa or more (TS. Gtoreq.1470 MPa), it is necessary to appropriately control the formation of the carbide.
For this reason, the holding time is required to be controlled to 20 to 1500 seconds at a holding temperature of 150 to 260 ℃.
If the holding temperature is lower than the lower limit of 150 ℃ or the holding time is short, the carbide distribution density in the phase-change phase becomes insufficient, and the delayed fracture resistance is deteriorated. On the other hand, at a high temperature of 260 ℃ higher than the upper limit of the holding temperature, coarsening of carbides in the crystal grains and at the bulk grain boundaries becomes remarkable, and there is a possibility that the delayed fracture resistance is deteriorated. In addition, if the holding time exceeds 1500 seconds, coarsening of carbides in the crystal grains and at the grain boundaries becomes remarkable, and there is a possibility that the delayed fracture resistance is deteriorated. Therefore, in the present invention, the continuous annealing is performed at a holding temperature of 150 to 260 ℃ for 20 to 1500 seconds.
From the viewpoint of stabilizing press formability such as adjustment of surface roughness and flattening of plate shape, the steel sheet thus obtained may be subjected to temper rolling. In this case, the temper rolling elongation is preferably set to 0.1% or more. The temper rolling elongation is preferably set to 0.6% or less. In this case, the surface finishing roll is a rough surface roll, and from the viewpoint of flattening the shape, the roughness Ra of the steel sheet is preferably adjusted to 0.8 μm or more. The roughness Ra of the steel sheet is preferably 1.8 μm or less.
As described above, the present steel sheet can be produced into a plated steel sheet by performing a hot-dip treatment during cooling after soaking in annealing, or by performing electroplating after continuous annealing. Examples of the type of the plating layer include Zn-based plating layers (Zn-based, zn-Ni-based, zn-Fe-based, etc.), and Al-based plating layers. In the case of hot dip plating, the steel sheet may be immersed in a plating bath during cooling from an annealing temperature to 420 ℃ at a first average cooling rate of 2 ℃/sec or more, and then heated to 480 to 600 ℃ to perform alloying treatment. After the alloying treatment, the alloy may be cooled at the second average cooling rate, and the holding treatment may be performed at a holding temperature of 150 to 260 ℃ for 20 to 1500 seconds.
After the hot dip plating and electroplating, heat treatment may be performed at a temperature of 260 ℃ or lower in order to reduce the intrusion of hydrogen into the steel.
As described above, according to the present invention, the delayed fracture resistance of the high-strength cold-rolled steel sheet is significantly improved, which contributes to improvement in the strength and weight reduction of the component by the application of the high-strength steel sheet. The thickness of the steel sheet of the present invention is preferably set to 0.5mm or more. The thickness of the sheet is preferably set to 2.0mm or less.
Next, the member and the method of manufacturing the same of the present invention will be described.
The member of the present invention is a member obtained by subjecting the steel sheet of the present invention to at least one of forming and joining processes. The method for producing a member according to the present invention includes a step of forming the steel sheet according to the present invention or a step of joining the steel sheet to produce a member.
The steel sheet of the present invention has a tensile strength of 1470MPa or more and excellent delayed fracture resistance. Therefore, the member obtained by using the steel sheet of the present invention is also high-strength, and has superior delayed fracture resistance characteristics compared to conventional high-strength members. In addition, if the member of the present invention is used, weight reduction can be achieved. Therefore, the member of the present invention can be suitably used for a vehicle body frame member, for example.
The molding may be performed by a general method such as press molding, without limitation. The joining process may be general welding such as spot welding and arc welding, rivet joining, and rivet joining, without limitation.
Examples
Hereinafter, examples of the present invention will be described.
After melting the steels having the compositions shown in table 1, they were cast into billets.
The billets were subjected to heat treatment and rolling as shown in Table 2, and steel sheets having a sheet thickness of 1.4mm were obtained.
Specifically, a billet having each component composition was heated at an average heating rate of 6 ℃/min to a heating holding temperature shown in table 2 of the billet surface thermometer, and the heating holding time shown in table 2 was held. Then, the billets were subjected to hot finish rolling at a finish rolling temperature of 870 ℃ for a residence time of 900 to 1000 ℃ shown in table 2, and were cooled at an average cooling rate of 50 ℃/sec in a range from the finish rolling temperature to 650 ℃.
Then, the hot-rolled steel sheet was cooled and coiled at a coiling temperature of 550 ℃, and the hot-rolled steel sheet was cold-rolled at a reduction ratio (cold rolling reduction ratio) of 50%, whereby a cold-rolled steel sheet was produced.
Then, the cold-rolled steel sheet was heated from 400 ℃ to an annealing temperature shown in table 2 at an average heating rate shown in table 2, and the soaking time shown in table 2 was soaked at the annealing temperature.
Then, the annealing was performed continuously by cooling from the annealing temperature (first cooling start temperature) to 420 ℃ (second cooling start temperature) at the first average cooling rate shown in table 2, further cooling from 420 ℃ (second cooling start temperature) to the cooling stop temperature shown in table 2 at the second average cooling rate shown in table 2, reheating as needed, and then holding at the holding temperature shown in table 2 for the holding time shown in table 2.
In addition, in the continuous annealing of No.12, the steel sheet was immersed in a hot dip galvanizing bath at 480 ℃ while being cooled to 420 ℃ at a first average cooling rate, and then was heated to 540 ℃ and held for 15 seconds to perform alloying treatment, thereby producing an alloyed hot dip galvanized steel sheet. Then, cooling was performed at the second average cooling rate shown in table 2, and holding treatment at the holding temperature and holding time shown in table 2 was performed.
Further, with regard to No.3, after continuous annealing, the obtained steel sheet was subjected to electroplating to obtain a Zn-plated steel sheet.
The obtained steel sheet was subjected to quantification of the metal structure by the above method, and further to tensile test and delayed fracture resistance evaluation test.
Specifically, the method for measuring the tissue is performed as follows.
The area ratios of martensite, bainite, and ferrite were measured as follows: after grinding the L-section (vertical section parallel to the rolling direction) of the steel sheet, the steel sheet was corroded with an ethanol nitrate solution, and 4 fields of view were observed with SEM at a magnification of 2000 times at a position 1/4 of the thickness from the surface of the steel sheet, and the photographed tissue photograph was subjected to image analysis to determine the area ratio of martensite, bainite, and ferrite. Here, martensite and bainite refer to structures that are gray or white in SEM. Here, bainite has the following characteristics. That is, the aspect ratio is 2.5 or more and the plate-like form is a slightly black structure compared with martensite. The width of the plate is 0.3-1.7 μm. The distribution density of carbide with diameter of 10-200 nm in bainite is 0-3/μm 2 . On the other hand, ferrite is a region of contrast that appears black in SEM. Note that, although a small amount of carbide, nitride, sulfide, or oxide is contained in the martensite or bainite, it is difficult to exclude them, and therefore the area ratio of the region containing them is defined as the area ratio.
The retained austenite (retained γ) is measured as follows: the surface layer of the steel sheet was subjected to chemical polishing with oxalic acid at 200 μm, and the retained austenite (retained γ) was obtained by an X-ray diffraction intensity method with respect to the sheet surface. Calculated from the integrated intensities of (200) α, (211) α, (220) α, (200) γ, (220) γ, and (311) γ diffraction plane peaks measured by Mo-kα rays.
The average grain size (primary γ grain size) of the prior austenite grains was measured as follows: after grinding the L-section (vertical section parallel to the rolling direction) of the steel sheet, the steel sheet was corroded with a reagent (for example, a saturated aqueous picric acid solution or a solution obtained by adding ferric chloride thereto) that corrodes the original γ grain boundaries, 4 fields of view were observed with an optical microscope at a magnification of 500 times at a position 1/4 the thickness from the steel sheet surface, and 15 lines were drawn at intervals of 10 μm or more in the actual length in the sheet thickness direction and the rolling direction in the obtained photographs, and the number of intersections of the grain boundaries and the lines was counted. The original γ particle diameter is obtained by multiplying a value obtained by dividing the bus length by the number of intersections by 1.13.
The number density A of precipitates having an equivalent circle diameter of 500nm or more was determined as follows: after grinding the L-section (vertical section parallel to the rolling direction) of the steel sheet, the steel sheet was continuously shot with SEM for 2mm in the region from 1/5 position to 4/5 position of the sheet thickness, i.e., the region from 1/5 position relative to the sheet thickness to 4/5 position from the center of the sheet thickness on the surface of the steel sheet 2 The number of such precipitates were measured from the SEM photograph, and the number density A of the precipitates having an equivalent circle diameter of 500nm or more was obtained. The magnification at which photographing was performed was 2000 times. In addition, when the composition analysis of each inclusion particle was performed, each inclusion particle was amplified 10000 times, and the above precipitate was analyzed. Here, the precipitate having an equivalent circle diameter of 500nm or more is Fe 23 (C,B) 6 And a precipitate containing B, and the presence or absence of a B peak was examined by elemental analysis using an energy dispersive X-ray spectrometry (EDS) with an acceleration voltage of 3kV, and the presence of a B peak was evaluated as the presence of the precipitate.
The tensile test is as follows: at a position 1/4 of the coil width, a JIS No. 5 tensile test piece was cut so that the direction of the roll was the longitudinal direction, and YP, TS, and El were evaluated by performing a tensile test (according to JIS Z2241).
The delayed fracture resistance was evaluated as follows.
A strip-shaped test piece having a rolling direction of 100mm at a right angle and a rolling direction of 30mm was cut from a position 1/4 of the coil width in the width direction of the steel sheet (coil) obtained. Cutting of the end face on the long side with a length of 100mmThe cutting was set to a shearing process, and in a state of the shearing process (without performing a mechanical process for removing burrs), a bending process was performed so that the burrs were on the outer circumferential side of the bend, and the test piece was fixed by a bolt while maintaining the shape of the test piece at the time of the bending. The clearance in the shearing process was set to 13% and the rake angle was set to 1 °. In the bending process, the tip bending radius was 10mm, and the angle inside the bending apex was set to 90 degrees (V bending). The punch used was a punch having a U-shape (the portion of the tip R is a semicircle and the thickness of the punch body is 2R) and having the same tip radius as the tip bending radius R, and the die used was a die having a corner R of 30 mm. The depth of the punch pressing the steel sheet was adjusted, and the steel sheet was formed so that the bending angle (angle inside the bending apex) of the tip was 90 degrees (V-shape). The test piece was fastened by a hydraulic jack so that the distance between the flange ends of the straight piece portions at the time of bending was the same as that at the time of bending (so that the opening of the straight piece portion due to springback was eliminated), and the bolt was fastened in this state. The bolts were fixed through holes of an oval shape (short axis 10mm, long axis 15 mm) provided in advance 10mm inward from the short side edges of the strip-shaped test piece. The obtained test piece after bolt tightening was immersed in a solution of 0.1 mass% ammonium thiocyanate in water and McIlvaine buffer at 1:1 and adjusting the pH to 8.0, and performing a delayed fracture resistance evaluation test. At this time, the temperature of the solution was set to 20℃per 1cm of the test piece 3 The liquid amount of the surface area was set to 20ml. After 24 hours, it was confirmed whether or not there was a visually identifiable level (length of 1mm or more) of cracking, and it was determined that the delayed fracture resistance was excellent when no cracking was observed.
The structure and properties of the obtained steel sheet are shown in table 3.
TABLE 3
Allowance area ratio (×1): total area ratio of bainite, ferrite, and retained austenite
(x2) primary gamma particle size: average grain size of prior austenite grains
(x3) a: number density of precipitates having equivalent circle diameter of 500nm or more
(x4) [ B ]: b content (mass%)
The steel sheet within the scope of the present invention has high strength and excellent delayed fracture resistance.
On the other hand, the C content of No.13 (steel M) is less than the lower limit of the prescribed value of the present invention, and TS is insufficient.
The C content of No.14 (steel N) exceeds the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
The Mn content of No.15 (steel O) is less than the lower limit of the prescribed value of the present invention, and the formation of martensite is insufficient, and the sufficient delayed fracture resistance is not obtained.
The P content of No.16 (steel P) exceeds the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
The S content of No.17 (steel Q) exceeds the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
The s omicron al content of No.18 (steel R) exceeds the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance was not obtained.
The N content of No.19 (steel S) exceeds the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
The content of Nb and Ti in No.20 (steel T) is smaller than the lower limit of the specified value of the present invention, and the primary gamma particle size is large, and the sufficient delayed fracture resistance is not obtained.
The amounts of Nb and Ti in No.21 (steel U) exceed the upper limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
The B content of No.22 (steel V) exceeds the upper limit of the prescribed value of the present invention, and coarse precipitates are large, and sufficient delayed fracture resistance is not obtained.
The B content of No.23 (steel W) is less than the lower limit of the prescribed value of the present invention, and sufficient delayed fracture resistance is not obtained.
No.24 (steel A) has a heating temperature (billet surface temperature (SRT)) lower than the lower limit of the specified value of the present invention, and the primary gamma particle size is large, and the sufficient delayed fracture resistance is not obtained.
The billet of No.25 (steel A) has a heating retention time of less than the lower limit of the predetermined value of the present invention, and the primary gamma particle size is large, and the sufficient delayed fracture resistance is not obtained.
The residence time at 900 to 1000℃of No.26 (steel A) exceeds the upper limit of the prescribed value of the present invention, and the number density A of the precipitate is excessive, and the sufficient delayed fracture resistance is not obtained.
The average heating rate at the time of annealing of No.27 (steel A) was less than the lower limit of the prescribed value of the present invention, and the primary gamma particle size was large, and sufficient delayed fracture resistance was not obtained.
The soaking time at the time of annealing of No.28 (steel A) exceeds the upper limit of the prescribed value of the present invention, and the primary gamma grain size is large, and sufficient delayed fracture resistance is not obtained.
The first average cooling rate at the time of annealing of No.29 (steel A) was less than the lower limit of the prescribed value of the present invention, and the formation of martensite was insufficient, and the sufficient delayed fracture resistance was not obtained.
The second average cooling rate at the time of annealing of No.30 (steel A) was less than the lower limit of the prescribed value of the present invention, the formation of martensite was insufficient, and the sufficient delayed fracture resistance was not obtained.
The cooling stop temperature at the time of annealing of No.31 (steel A) exceeds the upper limit of the prescribed value of the present invention, and the formation of martensite is insufficient, and the sufficient delayed fracture resistance is not obtained.
It is also known that: since the steel sheet of the present invention example has high strength and excellent delayed fracture resistance, the member obtained by forming the steel sheet of the present invention example and the member obtained by joining the steel sheet of the present invention example have high strength and excellent delayed fracture resistance as the steel sheet of the present invention example.
Claims (11)
1. A steel sheet comprising, in mass%, C:0.15% or more and 0.45% or less, si: less than 1.5%, mn: greater than 1.7%, P: less than 0.03%, S: less than 0.0040%, sol.al:0.20% or less, N: less than 0.005%, B:0.0015% or more and 0.0100% or less, and one or more of Nb and Ti: a total of 0.005% to 0.080%, and the balance of Fe and unavoidable impurities,
and has a structure in which the area ratio of martensite to the whole structure is 95% or more and 100% or less,
the average grain size of the prior austenite grains is less than 11.2 mu m,
the number density A of precipitates having an equivalent circle diameter of 500nm or more satisfies the following formula (1),
a (number/mm) 2 )≤8.5×10 5 ×[B]… (1)
Here, [ B ] represents the content (mass%) of B.
2. The steel sheet according to claim 1, further comprising, as the component composition, a composition selected from the group consisting of Cu:1.0% below and Ni:1.0% or less of one or two of the following.
3. The steel sheet according to claim 1 or 2, further comprising, as the component composition, a composition selected from the group consisting of Cr: less than 1.0%, mo: less than 0.3%, V: less than 0.5%, zr:0.2% below and W:0.2% or less of one or two or more kinds of the above-mentioned components.
4. The steel sheet according to any one of claims 1 to 3, further comprising, as the component composition, a composition selected from the group consisting of, in mass%, ca: less than 0.0030 percent, ce: less than 0.0030 percent, la: below 0.0030%, REM (excluding Ce, la): 0.0030% below and Mg:0.0030% or less of one or two or more kinds of the following.
5. The steel sheet according to any one of claims 1 to 4, further comprising, as the component composition, a composition selected from the group consisting of: 0.1% below and Sn:0.1% or less of one or two of the following.
6. The steel sheet according to any one of claims 1 to 5, wherein a plating layer is provided on the surface of the steel sheet.
7. A member comprising the steel sheet according to any one of claims 1 to 6.
8. A method for manufacturing a steel sheet, wherein,
heating a steel slab having the composition according to any one of claims 1 to 5 at an average heating rate of 10 ℃/min or less from 1000 ℃ to a heating holding temperature of 1250 ℃ or higher with a slab surface thermometer, and holding at the heating holding temperature for 30 minutes or more, and then,
the residence time at 900-1000 ℃ is set to 20-150 seconds, the finish rolling is performed under the condition that the finish rolling temperature is set to 850 ℃ or higher,
Cooling is performed at an average cooling rate of 40 ℃/sec or more in a range from the finish rolling temperature to 650 ℃,
then, coiling is performed at a coiling temperature of 650 ℃ or less, thereby producing a hot rolled steel sheet,
cold-rolling the hot-rolled steel sheet at a reduction of 40% or more to produce a cold-rolled steel sheet,
the following continuous anneals were performed:
setting the annealing temperature to 800-950 ℃, heating the cold-rolled steel sheet from 400 ℃ to the annealing temperature at an average heating rate of more than 1.0 ℃/s,
is kept at the annealing temperature for 600 seconds or less,
cooling from the annealing temperature to 420 ℃ at a first average cooling rate of above 2 ℃/sec,
cooling from 420 ℃ to a cooling stop temperature below 260 ℃ at a second average cooling rate of more than 10 ℃/sec,
then, the mixture is kept at a holding temperature of 150 to 260 ℃ for 20 to 1500 seconds.
9. The method for producing a steel sheet according to claim 8, wherein,
immersing the cold-rolled steel sheet in a plating bath while cooling at the first average cooling rate in the continuous annealing,
after dipping in the plating bath, the alloy is heated to 480-600 ℃.
10. The method for manufacturing a steel sheet according to claim 8, wherein the surface of the steel sheet is subjected to a plating treatment after the continuous annealing.
11. A method for producing a member, comprising the step of forming the steel sheet according to any one of claims 1 to 6, or joining the steel sheet to produce a member.
Applications Claiming Priority (3)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2021123190 | 2021-07-28 | ||
| JP2021-123190 | 2021-07-28 | ||
| PCT/JP2022/024964 WO2023008003A1 (en) | 2021-07-28 | 2022-06-22 | Steel sheet, member, method for producing said steel sheet, and method for producing said member |
Publications (1)
| Publication Number | Publication Date |
|---|---|
| CN117651786A true CN117651786A (en) | 2024-03-05 |
Family
ID=85087860
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CN202280050129.6A Pending CN117651786A (en) | 2021-07-28 | 2022-06-22 | Steel sheet, member, and method for producing same |
Country Status (7)
| Country | Link |
|---|---|
| US (1) | US20240318275A1 (en) |
| EP (1) | EP4350016A4 (en) |
| JP (1) | JP7226672B1 (en) |
| KR (1) | KR20240024947A (en) |
| CN (1) | CN117651786A (en) |
| MX (1) | MX2024001026A (en) |
| WO (1) | WO2023008003A1 (en) |
Families Citing this family (6)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| WO2024162176A1 (en) * | 2023-01-30 | 2024-08-08 | Ｊｆｅスチール株式会社 | Steel sheet, member, and production methods for these |
| WO2024162175A1 (en) * | 2023-01-30 | 2024-08-08 | Ｊｆｅスチール株式会社 | Steel sheet, member, and production methods for these |
| WO2024190764A1 (en) * | 2023-03-13 | 2024-09-19 | 日本製鉄株式会社 | Steel plate |
| WO2024190763A1 (en) * | 2023-03-13 | 2024-09-19 | 日本製鉄株式会社 | Steel sheet and method for producing same |
| WO2024203492A1 (en) * | 2023-03-28 | 2024-10-03 | Ｊｆｅスチール株式会社 | Plated steel sheet, member, and production methods for same |
| WO2024203491A1 (en) * | 2023-03-28 | 2024-10-03 | Ｊｆｅスチール株式会社 | Steel sheet, member, and methods for producing same |
Family Cites Families (8)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS4712882Y1 (en) | 1968-08-13 | 1972-05-12 | ||
| JPS60109606A (en) | 1983-11-16 | 1985-06-15 | Toyooki Kogyo Co Ltd | By-pass valve |
| JP4712882B2 (en) | 2008-07-11 | 2011-06-29 | 株式会社神戸製鋼所 | High strength cold-rolled steel sheet with excellent hydrogen embrittlement resistance and workability |
| EP3486346B1 (en) * | 2016-09-28 | 2020-08-12 | JFE Steel Corporation | Steel sheet and method of producing the same |
| MX2019002330A (en) | 2016-09-28 | 2019-07-04 | Jfe Steel Corp | Steel sheet and method for producing same. |
| CN113227415B (en) * | 2018-12-21 | 2023-05-05 | 杰富意钢铁株式会社 | Steel sheet, member, and method for producing same |
| US20220186339A1 (en) * | 2019-02-21 | 2022-06-16 | Jfe Steel Corporation | Hot-pressed member, cold-rolled steel sheet for hot pressing, and manufacturing methods therefor |
| US20220177992A1 (en) * | 2019-03-12 | 2022-06-09 | Jfe Steel Corporation | Hot press member, production method for steel sheet for hot press, and production method for hot press member |
-
2022
- 2022-06-22 KR KR1020247002167A patent/KR20240024947A/en unknown
- 2022-06-22 WO PCT/JP2022/024964 patent/WO2023008003A1/en active Application Filing
- 2022-06-22 JP JP2022559509A patent/JP7226672B1/en active Active
- 2022-06-22 US US18/579,914 patent/US20240318275A1/en active Pending
- 2022-06-22 CN CN202280050129.6A patent/CN117651786A/en active Pending
- 2022-06-22 EP EP22849078.5A patent/EP4350016A4/en active Pending
- 2022-06-22 MX MX2024001026A patent/MX2024001026A/en unknown
Also Published As
| Publication number | Publication date |
|---|---|
| EP4350016A1 (en) | 2024-04-10 |
| EP4350016A4 (en) | 2024-10-23 |
| US20240318275A1 (en) | 2024-09-26 |
| JPWO2023008003A1 (en) | 2023-02-02 |
| KR20240024947A (en) | 2024-02-26 |
| JP7226672B1 (en) | 2023-02-21 |
| WO2023008003A1 (en) | 2023-02-02 |
| MX2024001026A (en) | 2024-03-15 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| CN109642295B (en) | Steel sheet and method for producing same | |
| JP7226672B1 (en) | Steel plate, member and manufacturing method thereof | |
| US12077831B2 (en) | Steel sheet, member, and methods for producing them | |
| WO2018062380A1 (en) | Steel sheet and method for producing same | |
| EP2730666A1 (en) | Method for producing cold-rolled steel sheet | |
| CN111936657A (en) | High-strength steel sheet and method for producing same | |
| JP7140302B1 (en) | Steel plate, member and manufacturing method thereof | |
| CN113227415B (en) | Steel sheet, member, and method for producing same | |
| EP3901293B1 (en) | High-strength hot-dip galvanized steel sheet and manufacturing method therefor | |
| CN115244203B (en) | Hot rolled steel sheet | |
| CN114585759B (en) | High-strength steel sheet, impact absorbing member, and method for producing high-strength steel sheet | |
| CN114585758B (en) | High-strength steel sheet, impact absorbing member, and method for producing high-strength steel sheet | |
| EP4283006A1 (en) | Steel sheet, member, method for producing said steel sheet, and method for producing said member | |
| JP7226673B1 (en) | Steel plate, member and manufacturing method thereof | |
| JP7140301B1 (en) | Steel plate, member and manufacturing method thereof | |
| WO2024162176A1 (en) | Steel sheet, member, and production methods for these | |
| CN116897217A (en) | Steel sheet, member, and method for producing same |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| PB01 | Publication | ||
| PB01 | Publication | ||
| SE01 | Entry into force of request for substantive examination | ||
| SE01 | Entry into force of request for substantive examination |