WO2023032423A1 - High strength steel sheet, high strength plated steel sheet, methods for producing these, and member - Google Patents
High strength steel sheet, high strength plated steel sheet, methods for producing these, and member Download PDFInfo
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- WO2023032423A1 WO2023032423A1 PCT/JP2022/024762 JP2022024762W WO2023032423A1 WO 2023032423 A1 WO2023032423 A1 WO 2023032423A1 JP 2022024762 W JP2022024762 W JP 2022024762W WO 2023032423 A1 WO2023032423 A1 WO 2023032423A1
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 110
- 239000010959 steel Substances 0.000 title claims abstract description 110
- 238000000034 method Methods 0.000 title description 23
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 55
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 54
- 229910001563 bainite Inorganic materials 0.000 claims abstract description 20
- 238000004519 manufacturing process Methods 0.000 claims abstract description 18
- 239000000203 mixture Substances 0.000 claims abstract description 16
- 229910052757 nitrogen Inorganic materials 0.000 claims abstract description 12
- 239000013078 crystal Substances 0.000 claims abstract description 9
- 239000012535 impurity Substances 0.000 claims abstract description 6
- 238000010438 heat treatment Methods 0.000 claims description 59
- 238000001816 cooling Methods 0.000 claims description 21
- 238000005096 rolling process Methods 0.000 claims description 20
- 238000000137 annealing Methods 0.000 claims description 18
- 238000003303 reheating Methods 0.000 claims description 18
- 238000007747 plating Methods 0.000 claims description 10
- 238000005097 cold rolling Methods 0.000 claims description 4
- 229910052799 carbon Inorganic materials 0.000 abstract description 4
- 229910052758 niobium Inorganic materials 0.000 abstract description 4
- 229910052748 manganese Inorganic materials 0.000 abstract description 3
- 239000000470 constituent Substances 0.000 abstract description 2
- 229910052742 iron Inorganic materials 0.000 abstract description 2
- 229910052710 silicon Inorganic materials 0.000 abstract description 2
- 229910052698 phosphorus Inorganic materials 0.000 abstract 1
- 230000000694 effects Effects 0.000 description 43
- 230000007423 decrease Effects 0.000 description 15
- 238000005452 bending Methods 0.000 description 13
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- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 8
- 150000001247 metal acetylides Chemical class 0.000 description 8
- 229910001567 cementite Inorganic materials 0.000 description 7
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 7
- 230000000717 retained effect Effects 0.000 description 7
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- 238000005728 strengthening Methods 0.000 description 7
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- 229910000859 α-Fe Inorganic materials 0.000 description 6
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- 238000009864 tensile test Methods 0.000 description 4
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 3
- 238000004458 analytical method Methods 0.000 description 3
- 238000007796 conventional method Methods 0.000 description 3
- 238000005261 decarburization Methods 0.000 description 3
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- 229910052718 tin Inorganic materials 0.000 description 3
- 238000002441 X-ray diffraction Methods 0.000 description 2
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- 229910052797 bismuth Inorganic materials 0.000 description 2
- 238000005266 casting Methods 0.000 description 2
- 229910052804 chromium Inorganic materials 0.000 description 2
- 239000010960 cold rolled steel Substances 0.000 description 2
- 239000012141 concentrate Substances 0.000 description 2
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- 238000005336 cracking Methods 0.000 description 2
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- 229910052749 magnesium Inorganic materials 0.000 description 2
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- 229910001208 Crucible steel Inorganic materials 0.000 description 1
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
- 229910045601 alloy Inorganic materials 0.000 description 1
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- 229910052745 lead Inorganic materials 0.000 description 1
- 239000011159 matrix material Substances 0.000 description 1
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- 229910052750 molybdenum Inorganic materials 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 235000021110 pickles Nutrition 0.000 description 1
- 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 1
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- 229910052726 zirconium Inorganic materials 0.000 description 1
Images
Classifications
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- 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
- 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
- C21D8/0247—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
- 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
-
- 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
- C22C38/02—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
- C22C38/06—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
- C22C38/12—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
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
- C23C2/06—Zinc or cadmium or alloys based thereon
-
- C—CHEMISTRY; METALLURGY
- C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
- C25D—PROCESSES FOR THE ELECTROLYTIC OR ELECTROPHORETIC PRODUCTION OF COATINGS; ELECTROFORMING; APPARATUS THEREFOR
- C25D3/00—Electroplating: Baths therefor
- C25D3/02—Electroplating: Baths therefor from solutions
- C25D3/22—Electroplating: Baths therefor from solutions of zinc
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
Definitions
- the present invention relates to high-strength steel sheets, their manufacturing methods, and members.
- Patent Document 1 discloses a high-strength steel sheet with excellent workability and low-temperature toughness, and a method for manufacturing the same.
- Patent Document 2 discloses a high-strength steel sheet with excellent formability and impact resistance, and a method for producing a high-strength steel sheet with excellent formability and impact resistance.
- Patent Document 3 discloses a high-yield-ratio high-strength steel sheet and a method for producing the same.
- Patent Documents 1 and 2 do not consider the yield ratio.
- Patent Document 3 does not consider toughness.
- the present invention has been made in view of such circumstances, and an object of the present invention is to provide a high-strength steel sheet having a tensile strength of 1180 MPa or more and having both excellent bendability and toughness and a high yield ratio, and a method for producing the same.
- high strength means that the tensile strength TS measured according to JIS Z2241 is 1180 MPa or more.
- excellent bendability means that there is no cracking at the ridgeline of the bending apex in a bending test conducted in accordance with JISZ2248.
- excellent toughness means that the brittle-ductile transition temperature is -40°C or less in a Charpy impact test conducted in accordance with JIS Z2242.
- a high yield ratio means that the ratio YS/TS between yield strength and tensile strength measured according to JIS Z 2241 is 0.80 or more.
- the dislocations existing in the quenched martensite structure are mobile dislocations that easily cause slip motion at low stress, so the yield stress of the martensite structure is low.
- these dislocations migrate closer to the grain boundaries, where they become entangled and become immobile dislocations. Thereby, the yield ratio of the steel sheet can be increased.
- the present invention has been made based on the above findings. That is, the gist and configuration of the present invention are as follows.
- the component composition is further mass%, V: 0.100 or less, Mo: 0.500% or less, Cr: 1.00% or less, Cu: 1.00% or less, Ni: 0.50% or less, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.200% or less, W: 0.400% or less, Zr: 0.0200% or less, Ca: 0.0200% or less, Mg: 0.0200% or less, Co: 0.020% or less, REM: 0.0200% or less, Te: 0.020% or less,
- a steel slab having the chemical composition of [1] or [2] is hot-rolled to form a hot-rolled sheet, cold-rolling the hot-rolled sheet to obtain a cold-rolled sheet;
- the cold-rolled sheet is heated to a first heating temperature of 850° C. or higher and 920° C. or lower and held for 10 s or longer, and then heated to a second heating temperature of 1000° C. or higher and 1200° C. or lower at an average heating rate of 50° C./s or higher.
- An annealing step is performed in which the temperature is raised and cooled to 500 ° C.
- a method for producing a high-strength steel sheet wherein after the rolling step, a reheating step is performed in which the second cold-rolled sheet is held at a reheating temperature of 70° C. to 200° C. for 600 seconds or more to obtain a high-strength steel sheet.
- a high-strength steel sheet with a tensile strength of 1180 MPa or more, a high yield ratio, excellent bendability and toughness, and a method for producing the same.
- C 0.10% to 0.30% It has the effect of increasing the yield ratio. If the C content is less than 0.10%, the area ratios of martensite and bainite decrease, and a TS of 1180 MPa or more cannot be obtained. If the C content exceeds 0.30%, carbon borides of B and iron are formed during annealing, and a sufficient amount of B cannot be segregated at the prior austenite grain boundaries.
- the C content is preferably 0.11% or more. Also, the C content is preferably 0.28% or less.
- Si 0.20% or more and 1.20% or less
- Si is an element effective for solid solution strengthening and needs to be added in an amount of 0.20% or more.
- Si is an element that stabilizes ferrite and raises the transformation point. Therefore, if the Si content exceeds 1.20%, it becomes difficult to make the prior austenite grain size 10 ⁇ m or less.
- the Si content is preferably 0.50% or more.
- the Si content is preferably 1.10% or less.
- Mn 2.5% to 4.0% Mn is effective in improving hardenability. If the Mn content is less than 2.5%, the area ratios of martensite and bainite decrease, resulting in a decrease in strength. On the other hand, if the Mn content exceeds 4.0%, the segregation part becomes excessively hardened and the bendability decreases.
- the Mn content is preferably 2.8% or more.
- the Mn content is preferably 3.5% or less.
- P content is set to 0.050% or less because P segregates at prior austenite grain boundaries and lowers the toughness.
- the lower limit of the P content is not particularly set, and it may be 0%, but if it is less than 0.001%, the manufacturing cost increases, so 0.001% or more is preferable.
- the P content is preferably 0.025% or less.
- S 0.020% or less S segregates at prior austenite grain boundaries and lowers toughness, so the content is made 0.020% or less. There is no particular lower limit for the S content, but if it is less than 0.0001%, the manufacturing cost increases, so it is preferable to make it 0.0001% or more.
- the S content is preferably 0.018% or less.
- Al 0.10% or less
- Al is an element that acts as a deoxidizer, and in order to obtain such an effect, the Al content is preferably 0.005% or more.
- the Al content is preferably 0.05% or less.
- N 0.01% or less N forms nitrides with Nb and B, lowering the effect of adding Nb and B. Therefore, the content is set to 0.01% or less. Although the lower limit is not set, it is preferably 0.0001% or more from the viewpoint of manufacturing cost.
- Ti 0.100% or less Ti has the effect of fixing N in steel as TiN and suppressing the formation of BN and NbN. In order to obtain these effects, the Ti content is preferably 0.005% or more. On the other hand, if the Ti content exceeds 0.100%, coarse Ti carbides are formed on grain boundaries, resulting in a decrease in toughness. The Ti content is preferably 0.05% or less.
- Nb 0.002% or more and 0.050% or less Nb precipitates as a solid solution or as fine carbides and suppresses the growth of austenite grains during annealing.
- the Nb content is made 0.002% or more.
- the Nb content exceeds 0.050%, not only is the effect saturated, but also coarse Nb carbides are precipitated to lower the toughness.
- the Nb content is preferably 0.005% or more. Also, the Nb content is preferably 0.040% or less.
- B 0.0005% to 0.0050% B segregates at prior austenite grain boundaries and has the effect of increasing grain boundary strength. In order to obtain such an effect, the B content should be 0.0005% or more. On the other hand, if the B content exceeds 0.0050%, carbon borides are formed and the toughness decreases.
- the B content is preferably 0.0010% or more. Also, the B content is preferably 0.0030% or less.
- the high-strength cold-rolled steel sheet according to the present embodiment further has V: 0.100% or less, Mo: 0.500% or less, Cr: 1.00% or less, Cu: 1.00% or less, Ni: 0.50% or less, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.200% or less, W: 0.400% or less, Zr: 0.5% or less.
- 0200% or less Ca: 0.0200% or less, Mg: 0.0200% or less, Co: 0.020% or less, REM: 0.0200% or less, Te: 0.020% or less, Hf: 0.10% and Bi: 0.200% or less.
- V 0.100 or less V has the effect of forming fine carbides and increasing the strength. If the V content exceeds 0.100%, coarse V carbide precipitates and the toughness decreases.
- the lower limit of the V content is not particularly limited, and may be 0.000%, but is preferably 0.001% or more because it has the effect of forming fine carbides and increasing strength.
- Mo 0.500% or less Mo has the effect of improving hardenability and increasing the bainite and martensite fractions. If the Mo content exceeds 0.500%, the effect saturates. There is no particular lower limit to the Mo content, and it may be 0.000%. .
- Cr 1.00% or less Cr has the effect of improving hardenability and increasing the fractions of bainite and martensite. If the Cr content exceeds 1.00%, the effect saturates. There is no particular lower limit to the Cr content, and it may be 0.000%. .
- Cu 1.00% or less Cu has the effect of increasing the strength by solid solution. If the Cu content exceeds 1.00%, intergranular cracking is likely to occur. There is no particular lower limit to the Cu content, and it may be 0.000%, but it is preferably 0.01% or more because it has the effect of increasing the strength by solid solution.
- Ni 0.50% or less Ni has the effect of improving hardenability, but if the Ni content exceeds 0.50%, the effect saturates.
- the lower limit of the Ni content is not particularly limited, and may be 0.000%, but is preferably 0.01% or more because it has the effect of improving the hardenability.
- Sb 0.200% or less Sb has the effect of suppressing surface oxidation, nitriding and decarburization of the steel sheet, but if the Sb content exceeds 0.200%, the effect saturates.
- Sb content There is no particular lower limit for the Sb content, and it may be 0.000%, but it is preferably 0.001% or more because it has the effect of suppressing surface oxidation, nitriding, and decarburization of the steel sheet.
- Sn 0.200% or less Sn, like Sb, has the effect of suppressing surface oxidation, nitriding and decarburization of the steel sheet. If the Sn content exceeds 0.200%, the effect saturates. There is no particular lower limit to the Sn content, and it may be 0.000%.
- Ta 0.200% or less Ta has the effect of forming fine carbides and increasing the strength. If the Ta content exceeds 0.200%, coarse Ta carbides precipitate and the toughness decreases.
- the lower limit of the Ta content is not particularly limited, and it may be 0.000%.
- W 0.400% or less W has the effect of forming fine carbides and increasing the strength. If the W content exceeds 0.400%, coarse W carbide precipitates and the toughness decreases. There is no particular lower limit for the W content, and it may be 0.000%.
- Zr 0.0200% or less Zr has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the Zr content exceeds 0.0200%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Zr content, and it may be 0.000%. is preferred.
- Ca 0.0200% or less Ca can be used as a deoxidizer. If the Ca content exceeds 0.0200%, a large amount of Ca-based inclusions are formed, resulting in a decrease in toughness. There is no particular lower limit to the Ca content, and it may be 0.000%, but it is preferably 0.0001% or more because it can be used as a deoxidizer.
- Mg 0.0200% or less Mg can be used as a deoxidizer. If the Mg content exceeds 0.0200%, a large amount of Mg-based inclusions are formed, resulting in a decrease in toughness. There is no particular lower limit to the Mg content, and although it may be 0.000%, it is preferably 0.0001% or more because it can be used as a deoxidizer.
- Co 0.020% or less Co has the effect of increasing strength through solid-solution strengthening. If the Co content exceeds 0.020%, the effect saturates. There is no particular lower limit to the Co content, and it may be 0.000%.
- REM 0.0200% or less REM has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the REM content exceeds 0.0200%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit to the REM content, and it may be 0.000%. is preferred.
- Te 0.020% or less Te has the effect of making inclusions spherical to suppress stress concentration and improving toughness. If the Te content exceeds 0.020%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Te content, and it may be 0.000%. is preferred.
- Hf 0.10% or less Hf has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the Hf content exceeds 0.10%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Hf content, and it may be 0.000%. is preferred.
- Bi 0.200% or less Bi has the effect of reducing segregation and improving bendability. If the Bi content exceeds 0.200%, a large amount of inclusions is formed and the bendability is lowered. There is no particular lower limit to the Bi content, and it may be 0.000%.
- the balance other than the above components is Fe and unavoidable impurities. If the content of any of the above optional elements is less than the lower limit, the effect of the present invention will not be impaired.
- Martensite and Bainite Total Area Ratio of 95% or More Both martensite and bainite are hard phases and are necessary to achieve a TS of 1180 MPa or more. Therefore, the total area ratio of martensite and bainite should be 95% or more.
- the total area ratio of martensite and bainite is preferably 96% or more.
- the upper limit of the total area ratio of martensite and bainite is not particularly limited, and may be 100%.
- the steel structure may contain residual structures other than martensite and bainite.
- the residual structure includes ferrite, retained austenite and cementite.
- the total area ratio of the residual structure shall be 5% or less.
- the area ratio of each tissue is measured as follows.
- the area ratio of retained austenite was obtained by chemically polishing the rolled surface of a test piece taken from each steel plate to the position of 1/4t of the thickness of the steel plate, and examining the polished surface with an X-ray diffraction (XRD).
- XRD X-ray diffraction
- the diffraction intensity and the diffraction peak position are measured to calculate the volume fraction, and the number is taken as the area fraction of retained austenite.
- SEM images of 3 fields of view are taken at a magnification of 2000 for the observation surface.
- the total area ratio of martensite, bainite, and retained austenite and the area ratio of structures other than martensite, bainite, and retained austenite are determined by image analysis.
- the area ratio of retained austenite obtained by XRD is obtained.
- the average value of the three fields of view is taken as the area ratio of the tissue.
- Average crystal grain size of prior austenite grains 10 ⁇ m or less Toughness and bendability can be improved by refining the crystal grain size and complicating the crack growth path. Furthermore, it has the effect of increasing the yield strength by refining and strengthening the crystal grains. In order to obtain these effects, the average crystal grain size of prior austenite grains must be 10 ⁇ m or less.
- the average grain size of the prior austenite grains is preferably 9 ⁇ m or less.
- the lower limit of the average crystal grain size of the prior austenite grains is not particularly limited, it is preferably 1 ⁇ m or more from the viewpoint of production technology.
- the average grain size of prior austenite grains is measured as follows. After polishing the thickness cross-section of each steel sheet parallel to the rolling direction, it is corroded with picral, and the microstructure at the position of 1/4t of the thickness is photographed at a magnification of 2000 times for 3 fields of view SEM images. The grain size of each prior austenite grain is obtained by image analysis from the obtained structure image, and the average value of the three fields of view is taken as the average crystal grain size of the prior austenite grain.
- B concentration at prior austenite grain boundaries 0.10% or more by mass B segregates at prior austenite grain boundaries, thereby strengthening the grain boundaries and improving toughness and bendability. This effect can be obtained if the concentration of B in the prior austenite grain boundary is 0.10% or more in terms of mass %.
- the concentration of B in the prior austenite grain boundaries is preferably 0.15% or more, more preferably 0.20% or more in mass %.
- no upper limit is set for the B concentration at the prior austenite grain boundaries, it is preferably less than 20% in order to suitably prevent hard carbide borides from precipitating on the grain boundaries and further improve toughness.
- the B concentration at the prior austenite grain boundary is measured as follows.
- a needle-shaped sample is produced from a region containing prior austenite grain boundaries by an SEM-FIB (Focused Ion Beam) method.
- the obtained needle-like sample is subjected to 3DAP analysis using a three-dimensional atom probe (3DAP) device (LEAP4000XSi, manufactured by AMETEK). Measurements are performed in laser mode. From the number of B ions detected from the prior austenite grain boundary and the number of other ions, the B concentration at the prior austenite grain boundary is obtained.
- 3DAP three-dimensional atom probe
- FIG. 1 shows an example of a C-enriched region.
- FIG. 1A is a diagram showing observation results of C-enriched regions existing at block grain boundaries and packet grain boundaries.
- FIG. 1(B) is a diagram showing the observation results of C-enriched regions existing at prior austenite grain boundaries.
- the diagrams shown on the left are examples of the results of observation with a scanning transmission electron microscope (STEM), and a martensite grain boundary exists in the center of the diagram. I know there is.
- the figure shown on the right is an example of the observation result of the amount of C enrichment by STEM. From these figures, it can be seen that there are C-enriched regions along the martensite grain boundaries and extending to the base material sandwiching the martensite grain boundaries.
- C concentration in the C-enriched region 4.0 times or more the C content in the steel Grain boundary strength can be obtained. That is, the C concentration in the C-enriched region satisfies the following formula (2). C concentration in C-enriched region (mass%)/C content in steel (mass%) ⁇ 4.0 (2)
- the C concentration in the C-enriched region is preferably 4.5 times or more the C content in the steel.
- the C concentration is preferably 6% or less in order to suitably prevent cementite precipitation and suitably prevent a decrease in the dissolved C concentration.
- the C-enriched region Enriched width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary Strengthening can improve bendability and yield ratio. Therefore, the C-enriched region is present with an enrichment width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary. If the concentration width of the C concentration region is less than 3 nm, the above effect is small. On the other hand, if the width of the C-enriched region exceeds 100 nm, C cannot be sufficiently enriched at the grain boundary and near the grain boundary.
- the width of the C-enriched region is preferably 3.5 nm or more. Also, the width of the C-concentrated region is preferably 80 nm or less.
- the C-enriched region length of 100 nm or more in the direction parallel to the martensite grain boundary.
- the C-enriched region is made to exist with a length of 100 nm or more in the direction parallel to the martensite grain boundary. If the C-enriched region is less than 100 nm, breakage or yielding occurs from breaks in the C-enriched region.
- the C-enriched regions preferably have a length of 120 nm or more in the direction parallel to the martensite grain boundaries. There is no upper limit to the length of the C-enriched region along the martensite grain boundary, and the C-enriched region may exist so as to cover the entire length of the martensite grain boundary.
- the C concentration, concentration width and length of the C concentration area are measured as follows.
- a thin film sample is prepared by the SEM-FIB method from the region containing the martensite grain boundary, and the surface of C is analyzed by STEM and energy dispersive X-ray spectroscopy (EDS).
- EDS energy dispersive X-ray spectroscopy
- an analytical transmission electron microscope Talos F200X manufactured by FEI
- the thin film sample is tilted so that the martensite grain boundaries are parallel to the electron beam, and a surface analysis of a region of 200 ⁇ 500 nm is performed.
- the analysis length in the direction parallel to the martensite grain boundary (the direction along the martensite grain boundary) is set to 500 nm.
- a line profile with a length of 200 nm is obtained in the direction perpendicular to the martensite grain boundary by accumulating the surface analysis data in the direction parallel to the martensite grain boundary.
- the half value of the maximum value of the line profile is obtained, and the width of the line profile that is equal to or greater than the half value is defined as the thickened width of the C thickened region.
- EDS quantitative analysis is performed on the enriched width to quantify the C concentration in the C enriched region.
- the length of the C-enriched region is measured in the direction parallel to the martensite grain boundary, and the length is defined as the length of the C-enriched region along the martensite grain boundary.
- a high-strength steel sheet with a tensile strength of 1180 MPa or more can be provided.
- the tensile strength of the high-strength steel sheet is preferably 1250 MPa or more.
- the high-strength steel sheet described above may have a coating layer on at least one side.
- a coating layer any one of a hot-dip galvanized layer, an alloyed hot-dip galvanized layer, and an electro-galvanized layer is preferable.
- the composition of the plating layer is not particularly limited, and may be a known composition.
- the composition of the hot-dip galvanized layer is not particularly limited as long as it is a common one.
- the plating layer contains Fe: 20% by mass or less, Al: 0.001% by mass or more and 1.0% by mass or less, and further Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr , Co, Ca, Cu, Li, Ti, Be, Bi, and REM containing one or more selected from the group consisting of 0% by mass or more and 3.5% by mass or less in total, and the balance being Zn and unavoidable It has a composition consisting of organic impurities.
- the Fe content in the coating layer is less than 7% by mass, and in the case of an alloyed hot-dip galvanizing layer, in one example, the Fe content in the coating layer The amount is 7 mass % or more and 15 mass % or less, more preferably 8 mass % or more and 13 mass % or less.
- the coating weight is not particularly limited, but the coating weight per side of the high-strength steel sheet is preferably 20 to 80 g/m 2 .
- plating layers are formed on both the front and back surfaces of a high-strength steel sheet.
- a steel slab having the chemical composition described above is manufactured.
- a steel material is melted to obtain molten steel having the above composition.
- the smelting method is not particularly limited, and any of known smelting methods such as converter smelting and electric furnace smelting are suitable.
- the resulting molten steel is solidified to produce a steel slab (slab).
- the method for producing a steel slab from molten steel is not particularly limited, and a continuous casting method, an ingot casting method, a thin slab casting method, or the like can be used.
- the steel slab may be cooled once and then heated again before hot rolling, or the cast steel slab may be continuously hot rolled without being cooled to room temperature.
- the slab heating temperature is preferably 1100° C. or higher and preferably 1300° C. or lower in consideration of the rolling load and scale generation.
- the method of heating the slab is not particularly limited, but, for example, the slab can be heated in a conventional heating furnace.
- Hot rolling is not particularly limited and may be carried out according to a conventional method. Cooling after hot rolling is not particularly limited, and the steel is cooled to the coiling temperature. The hot-rolled sheet is then wound into a coil.
- the winding temperature is preferably 400° C. or higher. This is because, if the coiling temperature is 400° C. or higher, the coiling becomes easy without increasing the strength of the hot-rolled sheet. A winding temperature of 550° C. or higher is more preferable. Moreover, the coiling temperature is preferably 750.degree.
- the hot-rolled sheet may be heat-treated for the purpose of softening.
- the method of removing the scale is not particularly limited, but it is preferable to pickle the hot-rolled coil while unwinding it in order to remove the scale completely.
- the pickling method is not particularly limited, and a conventional method may be used.
- Cold rolling process After optionally washing the hot-rolled sheet from which scale has been removed, it is cold-rolled to obtain a cold-rolled sheet.
- the method of cold rolling is not particularly limited, and a conventional method may be followed.
- the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for 10 seconds or longer, and then to a second heating temperature of 1000°C or higher and 1200°C or lower at an average heating rate of 50°C/s or higher. Then, within 5 seconds after reaching the second heating temperature, the steel is cooled to 500°C or less at a cooling rate of 50°C/s or more.
- the cold-rolled sheet is heated to a first heating temperature of 850° C. or higher and 920° C. or lower and held for 10 seconds or longer.
- a first heating temperature 850° C. or higher and 920° C. or lower and held for 10 seconds or longer.
- annealing is performed at the first heating temperature in the austenite single-phase region. If the first heating temperature is lower than 850° C., ferrite is generated and the strength is lowered. On the other hand, if the first heating temperature exceeds 920° C., the austenite grain size exceeds 10 ⁇ m, and grain refinement cannot be achieved in subsequent steps, resulting in deterioration in bendability, toughness, and yield ratio.
- the first heating temperature is preferably 860° C. or higher. Also, the first heating temperature is preferably 900° C. or lower.
- Holding time at the first heating temperature 10 s or longer
- the holding time at the first heating temperature is 10 s or longer.
- the upper limit of the holding time at the first heating temperature is not particularly limited, but from the viewpoint of productivity, the holding time at the first heating temperature is preferably 60 seconds or less.
- the holding time at the first heating temperature is preferably 20 seconds or longer.
- Second heating temperature 1000° C. or more and 1200° C. or less After holding at the first heating temperature, annealing is performed at a high temperature while maintaining the austenite grain boundary at 10 ⁇ m or less, and a sufficient amount of B is segregated at the grain boundary. If the second heating temperature is less than 1000° C., diffusion of B is slow and grain boundary segregation is insufficient. If the second heating temperature exceeds 1200° C., the austenite grains grow rapidly and the austenite grain size exceeds 10 ⁇ m.
- the second heating temperature is preferably 1020° C. or higher.
- the second heating temperature is preferably 1150° C. or less.
- the average heating rate from the first heating temperature to the second heating temperature is 50°C/s or more.
- the austenite grain size grows to over 10 ⁇ m.
- the upper limit of the average heating rate from the first heating temperature to the second heating temperature is not particularly limited, it is preferably 120° C./s or less because excessive rapid heating is difficult to control.
- the average heating rate from the first heating temperature to the second heating temperature is preferably 80° C./s or higher.
- the average cooling rate from the second heating temperature to 500°C or lower shall be 50°C/s or higher. If the average cooling rate from the second heating temperature to 500° C. or lower is less than 50° C./s, grain growth occurs during cooling.
- the upper limit of the average cooling rate from the second heating temperature to 500° C. or lower is not particularly limited, it is preferably 120° C./s or lower in order to facilitate control.
- the average cooling rate from the second heating temperature to 500°C or lower is preferably 80°C/s or higher.
- Cooling stop temperature 500°C or less
- rapid cooling is performed to a cooling stop temperature of 500°C or less.
- the cooling stop temperature is preferably 450°C or less.
- the lower limit of the cooling stop temperature is not particularly limited, it is preferably 100° C. or higher.
- a plating process may be performed in which at least one side of the high-strength steel sheet is plated to obtain a high-strength plated steel sheet.
- the high-strength plated steel sheet may be heat-treated to alloy the coating layer of the high-strength plated steel sheet to obtain an alloyed plated steel sheet.
- a rolling step in which rolling is performed with an elongation of 0.5% or more
- the cold-rolled sheet is rolled at an elongation of 0.5% or more to obtain a second cold-rolled sheet , perform the rolling process.
- the cold-rolled sheet obtained through the steps up to this point contains many mobile dislocations. In this rolling process, mobile dislocations accumulate at grain boundaries and become entangled to become immobile dislocations. If the elongation rate is less than 0.5%, the effect is small.
- the elongation rate in the rolling process is preferably 0.6% or more.
- the upper limit of the elongation rate in the rolling process is not particularly set, it is preferably 2% or less, for example, in order to further reduce the load on the equipment.
- a reheating step in which the second cold-rolled sheet is held at a reheating temperature of 70°C or higher and 200°C or lower for 600 seconds or longer.
- the second cold-rolled sheet is tempered at a low temperature to form clusters. If the reheating temperature is less than 70° C., diffusion of C is slow and C does not concentrate near the grain boundary to a sufficient amount. On the other hand, if the reheating temperature exceeds 200° C., the tempering proceeds excessively and cementite precipitates. The cementite precipitated at the grain boundary tends to become a fracture initiation point, and the C concentration in the matrix around the cementite decreases, resulting in a decrease in bendability and toughness.
- the reheating temperature is preferably 90°C or higher. Also, the reheating temperature is preferably 190° C. or lower.
- Holding time at reheating temperature 600 s or more If the holding time at reheating temperature is less than 600 s, diffusion of C is slow and a sufficient amount of C concentration cannot be obtained.
- the upper limit of the holding time at the reheating temperature is not particularly limited, it is preferably 43200 seconds (0.5 days) or less in order to prevent precipitation of cementite.
- the holding time at the reheating temperature is preferably 800 s or more.
- the manufacturing conditions other than the above conditions can be according to the usual method.
- the high-strength steel sheet or high-strength galvanized steel sheet described above can be formed into a desired shape by press working and used as an automobile part.
- the automobile part may contain a steel sheet other than the high-strength steel sheet or the high-strength plated steel sheet according to the present embodiment as a material.
- the present high-strength steel sheet or high-strength galvanized steel sheet can be suitably used in general members used as frame structural parts or reinforcing parts, among automobile parts.
- a steel having the chemical composition shown in Table 1 and the balance being Fe and unavoidable impurities was melted in a converter to obtain a steel slab.
- the obtained slab was reheated, hot rolled, and coiled to obtain a hot rolled coil.
- the hot-rolled coil was pickled while being unwound, and cold-rolled.
- the thickness of the hot-rolled sheet was 3.0 mm, and the thickness of the cold-rolled sheet was 1.2 mm.
- Annealing was performed in a continuous hot-dip galvanizing line under the conditions shown in Table 2 to obtain a cold-rolled steel sheet, a hot-dip galvanized steel sheet (GI), and an alloyed hot-dip galvanized steel sheet (GA).
- the hot-dip galvanized steel sheet was immersed in a 460° C. plating bath to obtain a coating weight of 35 g/m 2 per side.
- the alloyed hot-dip galvanized steel sheet was manufactured by performing an alloying treatment at 520° C. for 40 seconds after adjusting the coating weight to 45 g/m 2 per side.
- the obtained steel sheets were subjected to rolling and reheating under the conditions shown in Table 2.
- the total area ratio of martensite and bainite, the prior austenite grain size, the B concentration of the prior austenite grain boundary, the C concentration (mass%) of the C enriched region at the martensite grain boundary / The C content (mass%) in the steel, the enriched width of the C-enriched region, and the length of the C-enriched region along the martensite grain boundary were evaluated.
- tensile strength, yield ratio, toughness and bendability were evaluated according to the methods described later. Table 3 shows the results.
- the Charpy impact test conforms to JIS Z. 2242. From the obtained steel sheet, a 90° A test piece with a V-notch of was taken. After that, a Charpy impact test was performed in a test temperature range of -120 to +120°C. A transition curve was obtained from the obtained brittle fracture surface ratio, and the temperature at which the brittle fracture surface ratio was 50% was determined as the brittle-ductile transition temperature. In addition, when the brittle-ductile transition temperature obtained from the Charpy test was -40°C or less, the toughness was judged to be good. In the table, when the brittle-ductile transition temperature is ⁇ 40° C. or less, the toughness is “excellent”, and when the brittle-ductile transition temperature is over ⁇ 40° C., the toughness is “poor”.
- the bending test is JIS Z. 2248.
- a strip-shaped test piece having a width of 30 mm and a length of 100 mm was taken from the obtained steel sheet so that the direction parallel to the rolling direction of the steel sheet was the axial direction of the bending test.
- a 90° V bending test was performed under the conditions of an indentation load of 100 kN and a pressing holding time of 5 seconds.
- bendability is evaluated by the pass rate of the bending test, and the maximum R at which the value R / t obtained by dividing the bending radius (R) by the plate thickness (t) is 5 or less (for example, when the plate thickness is 1.2 mm In the case, the bending radius is 7.0 mm), a bending test is performed on 5 samples, and then the presence or absence of crack generation at the ridge of the bending apex is evaluated. The bendability was judged to be good only in the case. In the table, the bendability is indicated as "excellent” only when the pass rate is 100%, and the bendability is indicated as "poor” in the other cases.
- the presence or absence of crack generation was evaluated by measuring the ridgeline portion of the bending apex with a digital microscope (RH-2000: manufactured by Hylox Co., Ltd.) at a magnification of 40 times.
- the TS of the present invention is 1180 MPa or more, the yield ratio is 0.80 or more, and the bendability and toughness are excellent.
- the comparative examples are inferior in at least one of TS, yield ratio, bendability and toughness.
Abstract
Description
C: 0.10%以上0.30以下,
Si: 0.20%以上1.20%以下,
Mn: 2.5%以上4.0%以下,
P: 0.050%以下,
S: 0.020%以下,
Al: 0.10%以下,
N: 0.01%以下,
Ti: 0.100%以下,
Nb: 0.002%以上0.050%以下及び
B: 0.0005%以上0.0050%以下
を含有し,残部がFe及び不可避的不純物からなり,下記式(1)を満足する成分組成を有し,
マルテンサイト及びベイナイトの面積率の合計が95%以上であり,
旧オーステナイト粒の平均結晶粒径が10μm以下であり,
旧オーステナイト粒界のB濃度が質量%で0.10%以上であり,
マルテンサイト粒界に沿ってC濃化領域を有し,
前記C濃化領域のC濃度が鋼中のC含有量の4.0倍以上であり,
前記マルテンサイト粒界と直交する方向において3nm以上100nm以下の濃化幅を有し,かつ前記マルテンサイト粒界に平行な方向において100nm以上の長さを有する,高強度鋼板。
([%N]/14)/([%Ti]/47.9)<1.0…(1)
式(1)において,[%N]及び[%Ti]はそれぞれN及びTiの鋼中含有量(質量%)を示す。 [1] in % by mass,
C: 0.10% or more and 0.30 or less,
Si: 0.20% or more and 1.20% or less,
Mn: 2.5% or more and 4.0% or less,
P: 0.050% or less,
S: 0.020% or less,
Al: 0.10% or less,
N: 0.01% or less,
Ti: 0.100% or less,
Nb: 0.002% or more and 0.050% or less and B: 0.0005% or more and 0.0050% or less, with the balance being Fe and unavoidable impurities, having a composition that satisfies the following formula (1): have
The total area ratio of martensite and bainite is 95% or more,
The average crystal grain size of the prior austenite grains is 10 μm or less,
The B concentration of the prior austenite grain boundary is 0.10% or more by mass%,
It has a C-enriched region along the martensite grain boundary,
The C concentration in the C-enriched region is 4.0 times or more the C content in the steel,
A high-strength steel sheet having a condensed width of 3 nm or more and 100 nm or less in a direction perpendicular to the martensite grain boundary and a length of 100 nm or more in a direction parallel to the martensite grain boundary.
([%N]/14)/([%Ti]/47.9)<1.0 (1)
In the formula (1), [%N] and [%Ti] indicate the content (% by mass) of N and Ti in the steel, respectively.
V: 0.100以下,
Mo: 0.500%以下,
Cr: 1.00%以下,
Cu: 1.00%以下,
Ni: 0.50%以下,
Sb: 0.200%以下,
Sn: 0.200%以下,
Ta: 0.200%以下,
W: 0.400%以下,
Zr: 0.0200%以下,
Ca: 0.0200%以下,
Mg: 0.0200%以下,
Co: 0.020%以下,
REM: 0.0200%以下,
Te: 0.020%以下,
Hf: 0.10%以下及び
Bi: 0.200%以下
のうちから選ばれる少なくとも1種の元素を含有する,前記[1]に記載の高強度鋼板。 [2] The component composition is further mass%,
V: 0.100 or less,
Mo: 0.500% or less,
Cr: 1.00% or less,
Cu: 1.00% or less,
Ni: 0.50% or less,
Sb: 0.200% or less,
Sn: 0.200% or less,
Ta: 0.200% or less,
W: 0.400% or less,
Zr: 0.0200% or less,
Ca: 0.0200% or less,
Mg: 0.0200% or less,
Co: 0.020% or less,
REM: 0.0200% or less,
Te: 0.020% or less,
The high-strength steel sheet according to [1] above, containing at least one element selected from Hf: 0.10% or less and Bi: 0.200% or less.
前記熱延板に冷間圧延を施して冷延板とし,
前記冷延板を,850℃以上920℃以下の第一加熱温度まで加熱して10s以上保持し,次いで,1000℃以上1200℃以下の第二加熱温度まで50℃/s以上の平均加熱速度で昇温し,該第二加熱温度に到達後5秒以内に,50℃/s以上の平均冷却速度で500℃以下まで冷却する,焼鈍工程を行い,
前記焼鈍工程の後,前記冷延板を伸長率0.5%以上にて圧延して第二冷延板を得る,圧延工程を行い,
前記圧延工程の後,前記第二冷延板を70℃以上200℃以下の再加熱温度に600s以上保持する再加熱工程を行なって高強度鋼板を得る,高強度鋼板の製造方法。 [4] A steel slab having the chemical composition of [1] or [2] is hot-rolled to form a hot-rolled sheet,
cold-rolling the hot-rolled sheet to obtain a cold-rolled sheet;
The cold-rolled sheet is heated to a first heating temperature of 850° C. or higher and 920° C. or lower and held for 10 s or longer, and then heated to a second heating temperature of 1000° C. or higher and 1200° C. or lower at an average heating rate of 50° C./s or higher. An annealing step is performed in which the temperature is raised and cooled to 500 ° C. or less at an average cooling rate of 50 ° C./s or more within 5 seconds after reaching the second heating temperature,
After the annealing step, performing a rolling step of rolling the cold-rolled sheet at an elongation of 0.5% or more to obtain a second cold-rolled sheet,
A method for producing a high-strength steel sheet, wherein after the rolling step, a reheating step is performed in which the second cold-rolled sheet is held at a reheating temperature of 70° C. to 200° C. for 600 seconds or more to obtain a high-strength steel sheet.
Cはマルテンサイト組織及びベイナイト組織を強化することに加え,旧オーステナイト粒界近傍に集積した転位に偏析して粒界を強化し,曲げ性,靭性及び降伏比を高める効果を有する。C含有量が0.10%未満では,マルテンサイト及びベイナイトの面積率が低下し,1180MPa以上のTSが得られない。C含有量が0.30%超では,焼鈍時にBと鉄との炭ホウ化物を形成し,旧オーステナイト粒界に十分な量のBを偏析させることができない。C含有量は,好ましくは0.11%以上とする。また,C含有量は,好ましくは0.28%以下とする。 C: 0.10% to 0.30% It has the effect of increasing the yield ratio. If the C content is less than 0.10%, the area ratios of martensite and bainite decrease, and a TS of 1180 MPa or more cannot be obtained. If the C content exceeds 0.30%, carbon borides of B and iron are formed during annealing, and a sufficient amount of B cannot be segregated at the prior austenite grain boundaries. The C content is preferably 0.11% or more. Also, the C content is preferably 0.28% or less.
Siは固溶強化に有効な元素であり,0.20%以上の添加を必要とする。一方,Siはフェライトを安定させる元素であり,変態点を上昇させる。そのため,Si含有量が1.20%超では,旧オーステナイト粒径を10μm以下とすることが困難となる。Si含有量は,好ましくは,0.50%以上とする。Si含有量は,好ましくは,1.10%以下とする。 Si: 0.20% or more and 1.20% or less Si is an element effective for solid solution strengthening and needs to be added in an amount of 0.20% or more. On the other hand, Si is an element that stabilizes ferrite and raises the transformation point. Therefore, if the Si content exceeds 1.20%, it becomes difficult to make the prior
Mnは焼入れ性向上に有効である。Mn含有量が2.5%未満では,マルテンサイト及びベイナイトの面積率が低下し,強度が低下する。一方で,Mn含有量が4.0%超では偏析部が過度に硬質化し曲げ性が低下する。Mn含有量は,好ましくは,2.8%以上とする。Mn含有量は,好ましくは,3.5%以下とする。 Mn: 2.5% to 4.0% Mn is effective in improving hardenability. If the Mn content is less than 2.5%, the area ratios of martensite and bainite decrease, resulting in a decrease in strength. On the other hand, if the Mn content exceeds 4.0%, the segregation part becomes excessively hardened and the bendability decreases. The Mn content is preferably 2.8% or more. The Mn content is preferably 3.5% or less.
Pは旧オーステナイト粒界に偏析し靭性を低下させるため,P含有量は0.050%以下とする。P含有量の下限は特に設けず,0%であってもよいが,0.001%未満とするのは製造コストを増加させるため,0.001%以上が好ましい。P含有量は,好ましくは0.025%以下とする。 P: 0.050% or less P content is set to 0.050% or less because P segregates at prior austenite grain boundaries and lowers the toughness. The lower limit of the P content is not particularly set, and it may be 0%, but if it is less than 0.001%, the manufacturing cost increases, so 0.001% or more is preferable. The P content is preferably 0.025% or less.
Sは旧オーステナイト粒界に偏析し靭性を低下させるため,0.020%以下とする。S含有量の下限は特に設けないが,0.0001%未満とするのは製造コストを増加させるため,0.0001%以上とすることが好ましい。S含有量は,好ましくは0.018%以下とする。 S: 0.020% or less S segregates at prior austenite grain boundaries and lowers toughness, so the content is made 0.020% or less. There is no particular lower limit for the S content, but if it is less than 0.0001%, the manufacturing cost increases, so it is preferable to make it 0.0001% or more. The S content is preferably 0.018% or less.
Alは脱酸材として作用する元素であり,そのような効果を得るためにAl含有量は0.005%以上とすることが好ましい。一方,Al含有量が0.10%超ではフェライトを生成しやすくなり強度が低下する。Al含有量は,0.05%以下とすることが好ましい。 Al: 0.10% or less Al is an element that acts as a deoxidizer, and in order to obtain such an effect, the Al content is preferably 0.005% or more. On the other hand, if the Al content exceeds 0.10%, ferrite tends to form and the strength decreases. The Al content is preferably 0.05% or less.
NはNb及びBと窒化物を形成し,Nb及びBの添加効果を下げる。そのため0.01%以下とする。下限は特に設けないが,製造コストの観点から0.0001%以上とすることが好ましい。 N: 0.01% or less N forms nitrides with Nb and B, lowering the effect of adding Nb and B. Therefore, the content is set to 0.01% or less. Although the lower limit is not set, it is preferably 0.0001% or more from the viewpoint of manufacturing cost.
Tiは鋼中のNをTiNとして固定し,BNやNbNの生成を抑制する効果を有する。これらの効果を得るために,Ti含有量は0.005%以上とすることが好ましい。一方,Ti含有量が0.100%超では,粗大なTi炭化物が粒界上に形成し,靭性が低下する。Ti含有量は好ましくは0.05%以下とする。 Ti: 0.100% or less Ti has the effect of fixing N in steel as TiN and suppressing the formation of BN and NbN. In order to obtain these effects, the Ti content is preferably 0.005% or more. On the other hand, if the Ti content exceeds 0.100%, coarse Ti carbides are formed on grain boundaries, resulting in a decrease in toughness. The Ti content is preferably 0.05% or less.
Nbは固溶もしくは微細な炭化物として析出し,オーステナイト粒の焼鈍中の成長を抑制する。このような効果を得るため,Nb含有量は0.002%以上とする。一方で,Nb含有量が0.050%超では効果が飽和するばかりか,粗大なNb炭化物が析出し靭性が低下する。Nb含有量は,好ましくは0.005%以上とする。また,Nb含有量は,好ましくは0.040%以下である。 Nb: 0.002% or more and 0.050% or less Nb precipitates as a solid solution or as fine carbides and suppresses the growth of austenite grains during annealing. In order to obtain such effects, the Nb content is made 0.002% or more. On the other hand, if the Nb content exceeds 0.050%, not only is the effect saturated, but also coarse Nb carbides are precipitated to lower the toughness. The Nb content is preferably 0.005% or more. Also, the Nb content is preferably 0.040% or less.
Bは旧オーステナイト粒界に偏析し粒界強度を高める効果を有する。そのような効果をえるためにB含有量は0.0005%以上とする。一方で,B含有量が0.0050%超では炭ホウ化物が形成し靭性が低下する。B含有量は,好ましくは0.0010%以上とする。また,B含有量は,好ましくは0.0030%以下とする。 B: 0.0005% to 0.0050% B segregates at prior austenite grain boundaries and has the effect of increasing grain boundary strength. In order to obtain such an effect, the B content should be 0.0005% or more. On the other hand, if the B content exceeds 0.0050%, carbon borides are formed and the toughness decreases. The B content is preferably 0.0010% or more. Also, the B content is preferably 0.0030% or less.
上述したB及びNbの添加効果を得るために,これらの元素と容易に結合するNはTiにより固定する必要がある。そのためにNのモル分率をTiのモル分率よりも小さくする。すなわち,上記式(1)を満足するようにN及びTiの鋼中含有量を調整する。なお,式(1)において,[%N]及び[%Ti]はそれぞれN及びTiの鋼中含有量(質量%)を示す。 ([%N]/14)/([%Ti]/47.9)<1.0 (1)
In order to obtain the effect of adding B and Nb described above, it is necessary to fix N, which easily bonds with these elements, by Ti. Therefore, the molar fraction of N is made smaller than the molar fraction of Ti. That is, the contents of N and Ti in the steel are adjusted so as to satisfy the above formula (1). In the formula (1), [%N] and [%Ti] indicate the content (% by mass) of N and Ti in the steel, respectively.
本実施形態に係る高強度冷延鋼板は,上記の成分組成に加えて,さらに質量%で,V:0.100以下,Mo:0.500%以下,Cr:1.00%以下,Cu:1.00%以下,Ni:0.50%以下,Sb:0.200%以下,Sn:0.200%以下,Ta:0.200%以下,W:0.400%以下,Zr:0.0200%以下,Ca:0.0200%以下,Mg:0.0200%以下,Co:0.020%以下,REM:0.0200%以下,Te:0.020%以下,Hf:0.10%以下及びBi:0.200%以下のうちから選ばれる少なくとも1種の元素を含有してもよい。 [Optional component]
In addition to the above chemical composition, the high-strength cold-rolled steel sheet according to the present embodiment further has V: 0.100% or less, Mo: 0.500% or less, Cr: 1.00% or less, Cu: 1.00% or less, Ni: 0.50% or less, Sb: 0.200% or less, Sn: 0.200% or less, Ta: 0.200% or less, W: 0.400% or less, Zr: 0.5% or less. 0200% or less, Ca: 0.0200% or less, Mg: 0.0200% or less, Co: 0.020% or less, REM: 0.0200% or less, Te: 0.020% or less, Hf: 0.10% and Bi: 0.200% or less.
Vは微細な炭化物を形成し強度を上げる効果を有する。V含有量が0.100%超では粗大なV炭化物が析出し靭性が低下する。V含有量の下限は特に限定されず,0.000%であってもよいが,微細な炭化物を形成し強度を上げる効果を有することから,0.001%以上とすることが好ましい。 V: 0.100 or less V has the effect of forming fine carbides and increasing the strength. If the V content exceeds 0.100%, coarse V carbide precipitates and the toughness decreases. The lower limit of the V content is not particularly limited, and may be 0.000%, but is preferably 0.001% or more because it has the effect of forming fine carbides and increasing strength.
Moは焼入れ性を向上しベイナイト及びマルテンサイト分率を高める効果を有する。Mo 含有量が0.500%超では効果が飽和する。Mo含有量の下限は特にされず,0.000%であってもよいが,焼入れ性を向上しベイナイト及びマルテンサイト分率を高める効果を有することから,0.010%以上とすることが好ましい。 Mo: 0.500% or less Mo has the effect of improving hardenability and increasing the bainite and martensite fractions. If the Mo content exceeds 0.500%, the effect saturates. There is no particular lower limit to the Mo content, and it may be 0.000%. .
Crは焼入れ性を向上しベイナイト及びマルテンサイト分率を高める効果を有する。Cr含有量が1.00%超では効果が飽和する。Cr含有量の下限は特にされず,0.000%であってもよいが,焼入れ性を向上しベイナイト及びマルテンサイト分率を高める効果を有することから,0.01%以上とすることが好ましい。 Cr: 1.00% or less Cr has the effect of improving hardenability and increasing the fractions of bainite and martensite. If the Cr content exceeds 1.00%, the effect saturates. There is no particular lower limit to the Cr content, and it may be 0.000%. .
Cuは固溶により強度を上昇する効果を有する。Cu含有量が1.00%超では粒界割れが生じやすくなる。Cu含有量の下限は特にされず,0.000%であってもよいが,固溶により強度を上昇する効果を有することから,0.01%以上とすることが好ましい。 Cu: 1.00% or less Cu has the effect of increasing the strength by solid solution. If the Cu content exceeds 1.00%, intergranular cracking is likely to occur. There is no particular lower limit to the Cu content, and it may be 0.000%, but it is preferably 0.01% or more because it has the effect of increasing the strength by solid solution.
Niは焼入れ性を向上する効果を有するが,Ni含有量が0.50%超では効果が飽和する。Ni含有量の下限は特にされず,0.000%であってもよいが,焼入れ性を向上する効果を有することから,0.01%以上とすることが好ましい。 Ni: 0.50% or less Ni has the effect of improving hardenability, but if the Ni content exceeds 0.50%, the effect saturates. The lower limit of the Ni content is not particularly limited, and may be 0.000%, but is preferably 0.01% or more because it has the effect of improving the hardenability.
Sbは鋼板の表面酸化や窒化,脱炭を抑制する効果を有するが,Sb含有量が0.200%超では効果が飽和する。Sb含有量の下限は特にされず,0.000%であってもよいが,鋼板の表面酸化や窒化,脱炭を抑制する効果を有することから,0.001%以上とすることが好ましい。 Sb: 0.200% or less Sb has the effect of suppressing surface oxidation, nitriding and decarburization of the steel sheet, but if the Sb content exceeds 0.200%, the effect saturates. There is no particular lower limit for the Sb content, and it may be 0.000%, but it is preferably 0.001% or more because it has the effect of suppressing surface oxidation, nitriding, and decarburization of the steel sheet.
SnはSbと同様に鋼板の表面酸化や窒化,脱炭を抑制する効果を有する。Sn含有量が0.200%超では効果が飽和する。Sn含有量の下限は特にされず,0.000%であってもよいが,鋼板の表面酸化や窒化,脱炭を抑制する効果を有することから,0.001%以上とすることが好ましい。 Sn: 0.200% or less Sn, like Sb, has the effect of suppressing surface oxidation, nitriding and decarburization of the steel sheet. If the Sn content exceeds 0.200%, the effect saturates. There is no particular lower limit to the Sn content, and it may be 0.000%.
Taは微細な炭化物を形成し強度を上げる効果を有する。Ta含有量が0.200%超では粗大なTa炭化物が析出し靭性が低下する。Ta含有量の下限は特にされず,0.000%であってもよいが,微細な炭化物を形成し強度を上げる効果を有することから,0.001%以上とすることが好ましい。 Ta: 0.200% or less Ta has the effect of forming fine carbides and increasing the strength. If the Ta content exceeds 0.200%, coarse Ta carbides precipitate and the toughness decreases. The lower limit of the Ta content is not particularly limited, and it may be 0.000%.
Wは微細な炭化物を形成し強度を上げる効果を有する。W含有量が0.400%超では,粗大なW炭化物が析出し靭性が低下する。W含有量の下限は特にされず,0.000%であってもよいが,微細な炭化物を形成し強度を上げる効果を有することから,0.001%以上とすることが好ましい。 W: 0.400% or less W has the effect of forming fine carbides and increasing the strength. If the W content exceeds 0.400%, coarse W carbide precipitates and the toughness decreases. There is no particular lower limit for the W content, and it may be 0.000%.
Zrは介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有する。Zr含有量が0.0200%超では介在物が多量に形成し靭性が低下する。Zr含有量の下限は特にされず,0.000%であってもよいが,介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有することから,0.0001%以上とすることが好ましい。 Zr: 0.0200% or less Zr has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the Zr content exceeds 0.0200%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Zr content, and it may be 0.000%. is preferred.
Caは脱酸材として用いることができる。Ca含有量が0.0200%超ではCa系介在物が多量に生成し靭性が低下する。Ca含有量の下限は特にされず,0.000%であってもよいが,脱酸材として用いることができることから,0.0001%以上とすることが好ましい。 Ca: 0.0200% or less Ca can be used as a deoxidizer. If the Ca content exceeds 0.0200%, a large amount of Ca-based inclusions are formed, resulting in a decrease in toughness. There is no particular lower limit to the Ca content, and it may be 0.000%, but it is preferably 0.0001% or more because it can be used as a deoxidizer.
Mgは脱酸材として用いることができる。Mg含有量が0.0200%超ではMg系介在物が多量に生成し靭性が低下する。Mg含有量の下限は特にされず,0.000%であってもよいが,脱酸材として用いることができることから,0.0001%以上とすることが好ましい。 Mg: 0.0200% or less Mg can be used as a deoxidizer. If the Mg content exceeds 0.0200%, a large amount of Mg-based inclusions are formed, resulting in a decrease in toughness. There is no particular lower limit to the Mg content, and although it may be 0.000%, it is preferably 0.0001% or more because it can be used as a deoxidizer.
Coは固溶強化で強度を上げる効果を有する。Co含有量が0.020%超では効果が飽和する。Co含有量の下限は特にされず,0.000%であってもよいが,固溶強化で強度を上げる効果を有することから,0.001%以上とすることが好ましい。 Co: 0.020% or less Co has the effect of increasing strength through solid-solution strengthening. If the Co content exceeds 0.020%, the effect saturates. There is no particular lower limit to the Co content, and it may be 0.000%.
REMは介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有する。REM含有量が0.0200%超では介在物が多量に形成し靭性が低下する。REM含有量の下限は特にされず,0.000%であってもよいが,介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有することから,0.0001%以上とすることが好ましい。 REM: 0.0200% or less REM has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the REM content exceeds 0.0200%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit to the REM content, and it may be 0.000%. is preferred.
Teは介在物の形状を球状化して応力集中を抑制し,靭性を向上させる効果を有する。Te含有量が0.020%超では介在物が多量に形成し靭性が低下する。Te含有量の下限は特にされず,0.000%であってもよいが,介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有することから,0.001%以上とすることが好ましい。 Te: 0.020% or less Te has the effect of making inclusions spherical to suppress stress concentration and improving toughness. If the Te content exceeds 0.020%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Te content, and it may be 0.000%. is preferred.
Hfは介在物の形状を球状化して応力集中を抑制し,靭性を向上させる効果を有する。Hf含有量が0.10%超では介在物が多量に形成し靭性が低下する。Hf含有量の下限は特にされず,0.000%であってもよいが,介在物の形状を球状化し応力集中を抑制し靭性を向上させる効果を有することから,0.01%以上とすることが好ましい。 Hf: 0.10% or less Hf has the effect of spheroidizing inclusions, suppressing stress concentration, and improving toughness. If the Hf content exceeds 0.10%, a large amount of inclusions is formed and the toughness is lowered. There is no particular lower limit for the Hf content, and it may be 0.000%. is preferred.
Biは偏析を軽減して曲げ性を向上させる効果を有する。Bi含有量が0.200%超では介在物が多量に形成し曲げ性が低下する。Bi含有量の下限は特にされず,0.000%であってもよいが,偏析を軽減し曲げ性を向上させる効果を有することから,0.001%以上とすることが好ましい。 Bi: 0.200% or less Bi has the effect of reducing segregation and improving bendability. If the Bi content exceeds 0.200%, a large amount of inclusions is formed and the bendability is lowered. There is no particular lower limit to the Bi content, and it may be 0.000%.
次に,高強度鋼板の鋼組織について説明する。 [Steel structure]
Next, the steel structure of high-strength steel sheets will be explained.
マルテンサイト及びベイナイトともに硬質相であり,1180MPa以上のTSを達成するために必要である。そのため,マルテンサイ及びベイナイトの面積率の合計は95%以上とする。マルテンサイ及びベイナイトの面積率の合計は,好ましくは96%以上である。マルテンサイ及びベイナイトの面積率の合計の上限は特に限定されず,100%であってもよい。 Martensite and Bainite: Total Area Ratio of 95% or More Both martensite and bainite are hard phases and are necessary to achieve a TS of 1180 MPa or more. Therefore, the total area ratio of martensite and bainite should be 95% or more. The total area ratio of martensite and bainite is preferably 96% or more. The upper limit of the total area ratio of martensite and bainite is not particularly limited, and may be 100%.
結晶粒径を微細化して亀裂進展経路を複雑化することで,靭性及び曲げ性の向上が可能である。さらに結晶粒を微細化して強化することにより降伏強度を上げる効果を有する。これら効果を得るためには,旧オーステナイト粒の平均結晶粒径を10μm以下とする必要がある。旧オーステナイト粒の平均結晶粒径は好ましくは9μm以下である。旧オーステナイト粒の平均結晶粒径の下限は特に限定されないが,生産技術上の観点から,1μm以上とすることが好ましい。 Average crystal grain size of prior austenite grains: 10 µm or less Toughness and bendability can be improved by refining the crystal grain size and complicating the crack growth path. Furthermore, it has the effect of increasing the yield strength by refining and strengthening the crystal grains. In order to obtain these effects, the average crystal grain size of prior austenite grains must be 10 μm or less. The average grain size of the prior austenite grains is preferably 9 μm or less. Although the lower limit of the average crystal grain size of the prior austenite grains is not particularly limited, it is preferably 1 μm or more from the viewpoint of production technology.
Bは旧オーステナイト粒界に偏析することで粒界を強化し,靭性及び曲げ性を向上させることができる。旧オーステナイト粒界のB濃度が質量%で0.10%以上であれば,該効果が得られる。旧オーステナイト粒界のB濃度は,好ましくは,質量%で0.15%以上,より好ましくは0.20%以上である。旧オーステナイト粒界のB濃度の上限は設けないが,硬質の炭ホウ化物が粒界上に析出することを好適に防ぎ,靭性をより向上させるため,このましくは20%未満である。 B concentration at prior austenite grain boundaries: 0.10% or more by mass B segregates at prior austenite grain boundaries, thereby strengthening the grain boundaries and improving toughness and bendability. This effect can be obtained if the concentration of B in the prior austenite grain boundary is 0.10% or more in terms of mass %. The concentration of B in the prior austenite grain boundaries is preferably 0.15% or more, more preferably 0.20% or more in mass %. Although no upper limit is set for the B concentration at the prior austenite grain boundaries, it is preferably less than 20% in order to suitably prevent hard carbide borides from precipitating on the grain boundaries and further improve toughness.
マルテンサイト粒界及びマルテンサイト粒界を挟む母相をC濃化により強化することで,曲げ性及び降伏比を向上することができる。なお,本明細書中において,「マルテンサイト粒界」は,マルテンサイト及びベイナイトに存在する旧オーステナイト粒界,ブロック粒界及びパケット粒界をいずれも含む。図1に,C濃化領域の一例を示す。図1(A)は,ブロック粒界及びパケット粒界に存在するC濃化領域の観察結果を示す図である。図1(B)は,旧オーステナイト粒界に存在するC濃化領域の観察結果を示す図である。図1(A),(B)において,それぞれ左に示す図は,走査透過電子顕微鏡(Scanning Transmission Electron Microscope:STEM)での観察結果の一例であり,図中央にマルテンサイト粒界が存在していることがわかる。右に示す図は,STEMでのC濃化量の観察結果の一例である。これら図から,マルテンサイト粒界に沿って,マルテンサイト粒界を挟む母材にかけてC濃化領域が存在していることがわかる。 C-Enriched Region Bendability and yield ratio can be improved by strengthening the martensite grain boundary and the parent phase sandwiching the martensite grain boundary by C-enrichment. In this specification, "martensite grain boundaries" include prior austenite grain boundaries, block grain boundaries and packet grain boundaries existing in martensite and bainite. FIG. 1 shows an example of a C-enriched region. FIG. 1A is a diagram showing observation results of C-enriched regions existing at block grain boundaries and packet grain boundaries. FIG. 1(B) is a diagram showing the observation results of C-enriched regions existing at prior austenite grain boundaries. In FIGS. 1(A) and (B), the diagrams shown on the left are examples of the results of observation with a scanning transmission electron microscope (STEM), and a martensite grain boundary exists in the center of the diagram. I know there is. The figure shown on the right is an example of the observation result of the amount of C enrichment by STEM. From these figures, it can be seen that there are C-enriched regions along the martensite grain boundaries and extending to the base material sandwiching the martensite grain boundaries.
C濃化領域において,鋼中のC含有量の4.0倍以上にCが濃化することで,十分な結晶粒界強度を得ることができる。すなわち,C濃化領域のC濃度は,下記式(2)を満たす。
C濃化領域のC濃度(質量%)/鋼中のC含有量(質量%)≧4.0…(2)
C濃化領域のC濃度は,好ましくは鋼中のC含有量の4.5倍以上である。C濃化領域のC濃度の上限は特に設けないが,セメンタイトの析出を好適に防いで固溶C濃度の低下を好適に防ぐため,C濃度が6%以下であることが好ましい。 C concentration in the C-enriched region: 4.0 times or more the C content in the steel Grain boundary strength can be obtained. That is, the C concentration in the C-enriched region satisfies the following formula (2).
C concentration in C-enriched region (mass%)/C content in steel (mass%) ≥ 4.0 (2)
The C concentration in the C-enriched region is preferably 4.5 times or more the C content in the steel. Although there is no particular upper limit for the C concentration in the C-enriched region, the C concentration is preferably 6% or less in order to suitably prevent cementite precipitation and suitably prevent a decrease in the dissolved C concentration.
図1に示すように,マルテンサイト粒界だけではなく,マルテンサイト粒界を挟む母相までC濃化により強化することで,曲げ性及び降伏比の向上が可能である。よって,C濃化領域を,マルテンサイト粒界に直交する方向において3nm以上100nm以下の濃化幅で存在させる。C濃化領域の濃化幅が3nm未満では上記効果が小さい。一方で,C濃化領域の幅が100nm超ではCを粒界及び粒界近傍に十分に濃化させることができない。C濃化領域の幅は,好ましくは3.5nm以上である。また,C濃化領域の幅は,好ましくは80nm以下である。 C-enriched region: Enriched width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary Strengthening can improve bendability and yield ratio. Therefore, the C-enriched region is present with an enrichment width of 3 nm or more and 100 nm or less in the direction perpendicular to the martensite grain boundary. If the concentration width of the C concentration region is less than 3 nm, the above effect is small. On the other hand, if the width of the C-enriched region exceeds 100 nm, C cannot be sufficiently enriched at the grain boundary and near the grain boundary. The width of the C-enriched region is preferably 3.5 nm or more. Also, the width of the C-concentrated region is preferably 80 nm or less.
優れた曲げ性及び降伏比を得るために,C偏析によりマルテンサイト粒界をネットワーク状に強化することが重要である。そのため,C濃化領域をマルテンサイト粒界に平行な方向において100nm以上の長さ存在させる。C濃化領域が100nm未満では,C濃化領域の切れ目から破壊や降伏が生じる。C濃化領域は,好ましくはマルテンサイト粒界に平行な方向において120nm以上の長さ存在する。C濃化領域のマルテンサイト粒界に沿った長さの上限はなく,マルテンサイト粒界の全長を覆うようにC濃化領域が存在してもよい。 C-enriched region: length of 100 nm or more in the direction parallel to the martensite grain boundary In order to obtain excellent bendability and yield ratio, it is important to strengthen the martensite grain boundary in a network by C segregation. Therefore, the C-enriched region is made to exist with a length of 100 nm or more in the direction parallel to the martensite grain boundary. If the C-enriched region is less than 100 nm, breakage or yielding occurs from breaks in the C-enriched region. The C-enriched regions preferably have a length of 120 nm or more in the direction parallel to the martensite grain boundaries. There is no upper limit to the length of the C-enriched region along the martensite grain boundary, and the C-enriched region may exist so as to cover the entire length of the martensite grain boundary.
次いで,スラブ加熱された鋼スラブを熱間圧延して熱延板とする。熱間圧延は特に制限はなく,常法に従い行えばよい。熱間圧延後の冷却もとくに制限はなく,巻取り温度まで冷却する。次いで,熱延板をコイルに巻取る。巻取り温度は400℃以上とすることが好ましい。巻取り温度が400℃以上であれば,熱延板の強度が上昇することなく巻き取りが容易となるからである。巻取り温度は550℃以上がより好ましい。また,スケールが厚く生成することを好適に防ぎ,歩留りをより向上するために,巻取り温度は750℃以下とすることが好ましい。なお,酸洗前に,軟質化を目的として熱延板に熱処理を行ってもよい。 [Hot rolling process]
Then, the slab-heated steel slab is hot-rolled into a hot-rolled sheet. Hot rolling is not particularly limited and may be carried out according to a conventional method. Cooling after hot rolling is not particularly limited, and the steel is cooled to the coiling temperature. The hot-rolled sheet is then wound into a coil. The winding temperature is preferably 400° C. or higher. This is because, if the coiling temperature is 400° C. or higher, the coiling becomes easy without increasing the strength of the hot-rolled sheet. A winding temperature of 550° C. or higher is more preferable. Moreover, the coiling temperature is preferably 750.degree. In addition, before pickling, the hot-rolled sheet may be heat-treated for the purpose of softening.
任意で,コイルに巻取った熱延板のスケールを除去する。スケールを除去する方法は特に限定されないが,スケール完全に除去するために,熱延コイルを巻戻しながら酸洗を行うことが好ましい。酸洗方法はとくに限定されず,常法に従えばよい。 [Pickling process]
Optionally, descale the hot-rolled sheet wound into coils. The method of removing the scale is not particularly limited, but it is preferable to pickle the hot-rolled coil while unwinding it in order to remove the scale completely. The pickling method is not particularly limited, and a conventional method may be used.
任意でスケールを除去した熱延板を適宜洗浄した後,冷間圧延して冷延板とする。冷間圧延の方法は特に限定されず常法に従えばよい。 [Cold rolling process]
After optionally washing the hot-rolled sheet from which scale has been removed, it is cold-rolled to obtain a cold-rolled sheet. The method of cold rolling is not particularly limited, and a conventional method may be followed.
次いで,冷延板を,850℃以上920℃以下の第一加熱温度まで加熱して10s以上保持し,次いで,1000℃以上1200℃以下の第二加熱温度まで50℃/s以上の平均加熱速度で昇温し,該第二加熱温度に到達後5秒以内に,冷却速度50℃/s以上の冷却速度で500℃以下まで冷却する,焼鈍工程を行う。 [Annealing process]
Next, the cold-rolled sheet is heated to a first heating temperature of 850°C or higher and 920°C or lower and held for 10 seconds or longer, and then to a second heating temperature of 1000°C or higher and 1200°C or lower at an average heating rate of 50°C/s or higher. Then, within 5 seconds after reaching the second heating temperature, the steel is cooled to 500°C or less at a cooling rate of 50°C/s or more.
次いで,冷延板を,850℃以上920℃以下の第一加熱温度まで加熱して10s以上保持する。マルテンサイト及びベイナイト主体の組織を得るため,オーステナイト単相域の第一加熱温度にて焼鈍を行う。第一加熱温度が850℃未満では,フェライトが生成し強度が低下する。一方,第一加熱温度が920℃超では,オーステナイト粒径が10μmを超え,以降の工程では細粒化できないため,曲げ性及び靭性及び降伏比が低下する。第一加熱温度は,好ましくは860℃以上である。また,第一加熱温度は,好ましくは900℃以下である。 First heating temperature of 850° C. or higher and 920° C. or lower Next, the cold-rolled sheet is heated to a first heating temperature of 850° C. or higher and 920° C. or lower and held for 10 seconds or longer. In order to obtain a structure mainly composed of martensite and bainite, annealing is performed at the first heating temperature in the austenite single-phase region. If the first heating temperature is lower than 850° C., ferrite is generated and the strength is lowered. On the other hand, if the first heating temperature exceeds 920° C., the austenite grain size exceeds 10 μm, and grain refinement cannot be achieved in subsequent steps, resulting in deterioration in bendability, toughness, and yield ratio. The first heating temperature is preferably 860° C. or higher. Also, the first heating temperature is preferably 900° C. or lower.
第一加熱温度における保持時間は,10s以上とする。第一加熱温度にて10s以上保持することで,オーステナイト粒径の成長と,Nb炭化物によるピン止めもしくは固溶による成長抑制とが釣り合う。保持時間が10s未満ではオーステナイト粒が成長途中であり,続く急速加熱中にNb炭化物によるピン止めもしくは固溶による成長抑制の効果が発現せず,旧オーステナイト粒径が10μmを超える。第一加熱温度における保持時間の上限は特に限定しないが,生産性の観点から,第一加熱温度における保持時間は60s以下とすることが好ましい。第一加熱温度における保持時間は,好ましくは20s以上である。 Holding time at the first heating temperature: 10 s or longer The holding time at the first heating temperature is 10 s or longer. By maintaining the first heating temperature for 10 seconds or more, the growth of the austenite grain size is balanced with the suppression of the growth by pinning or solid solution by the Nb carbide. If the holding time is less than 10 s, the austenite grains are in the process of growing, and during the subsequent rapid heating, the effect of pinning by Nb carbide or growth inhibition by solid solution does not appear, and the prior austenite grain size exceeds 10 μm. The upper limit of the holding time at the first heating temperature is not particularly limited, but from the viewpoint of productivity, the holding time at the first heating temperature is preferably 60 seconds or less. The holding time at the first heating temperature is preferably 20 seconds or longer.
第一加熱温度での保持後,オーステナイト粒界を10μm以下に維持したまま高温で焼鈍し,Bを十分な量粒界偏析させる。第二加熱温度が1000℃未満ではBの拡散が遅く,粒界偏析が不十分である。第二加熱温度が1200℃超ではオーステナイト粒の成長が早く,オーステナイト粒径が10μm超となる。第二加熱温度は,好ましくは1020℃以上とする。第二加熱温度は,好ましくは1150℃以下である。 Second heating temperature of 1000° C. or more and 1200° C. or less After holding at the first heating temperature, annealing is performed at a high temperature while maintaining the austenite grain boundary at 10 μm or less, and a sufficient amount of B is segregated at the grain boundary. If the second heating temperature is less than 1000° C., diffusion of B is slow and grain boundary segregation is insufficient. If the second heating temperature exceeds 1200° C., the austenite grains grow rapidly and the austenite grain size exceeds 10 μm. The second heating temperature is preferably 1020° C. or higher. The second heating temperature is preferably 1150° C. or less.
第一加熱温度から第二加熱温度までの平均加熱速度は50℃/s以上とする。第一加熱温度から第二加熱温度までの平均加熱速度が50℃/s未満では,オーステナイト粒径が10μm超まで成長する。第一加熱温度から第二加熱温度までの平均加熱速度の上限は特に限定されないが,過度の急速加熱は制御が困難であるため,好ましくは120℃/s以下とする。第一加熱温度から第二加熱温度までの平均加熱速度は,好ましくは80℃/s以上である。 Average heating rate: 50°C/s or more The average heating rate from the first heating temperature to the second heating temperature is 50°C/s or more. When the average heating rate from the first heating temperature to the second heating temperature is less than 50° C./s, the austenite grain size grows to over 10 μm. Although the upper limit of the average heating rate from the first heating temperature to the second heating temperature is not particularly limited, it is preferably 120° C./s or less because excessive rapid heating is difficult to control. The average heating rate from the first heating temperature to the second heating temperature is preferably 80° C./s or higher.
第二加熱温度まで到達後は,該第二加熱温度にて保持することなく,第二加熱温度に到達後5秒以内に急冷を開始し,50℃/s以上の平均冷却速度にて500℃以下まで急冷を行なう。これにより,オーステナイト粒径10μm以下でBが0.1%以上粒界偏析した鋼組織を得ることができる。第二加熱温度にて保持すると,粒成長が速やかに始まるため,第二加熱温度に到達後ただちに冷却を開始する。 Within 5 seconds after reaching the second heating temperature, cool to 500 ° C or less at an average cooling rate of 50 ° C/s or more. Rapid cooling is started within 5 seconds after reaching the heating temperature, and is rapidly cooled to 500° C. or less at an average cooling rate of 50° C./s or more. As a result, a steel structure having an austenite grain size of 10 μm or less and B segregating at grain boundaries of 0.1% or more can be obtained. When the second heating temperature is maintained, grain growth starts quickly, so cooling is started immediately after reaching the second heating temperature.
第二加熱温度に到達後の冷却において,第二加熱温度から500℃以下までの平均冷却速度は50℃/s以上とする。第二加熱温度から500℃以下までの平均冷却速度が50℃/s未満では,冷却中に粒成長が生じる。第二加熱温度から500℃以下までの平均冷却速度の上限は特に限定されないが,制御を容易とするため,好ましくは120℃/s以下とする。第二加熱温度から500℃以下までの平均冷却速度は,好ましくは80℃/s以上とする。 Average cooling rate: 50°C/s or higher In the cooling after reaching the second heating temperature, the average cooling rate from the second heating temperature to 500°C or lower shall be 50°C/s or higher. If the average cooling rate from the second heating temperature to 500° C. or lower is less than 50° C./s, grain growth occurs during cooling. Although the upper limit of the average cooling rate from the second heating temperature to 500° C. or lower is not particularly limited, it is preferably 120° C./s or lower in order to facilitate control. The average cooling rate from the second heating temperature to 500°C or lower is preferably 80°C/s or higher.
また,フェライト変態を抑制するため,500℃以下の冷却停止温度まで急冷を行う。冷却停止温度は,好ましくは450℃以下とする。冷却停止温度の下限は特に限定されないが,100℃以上とすることが好ましい。 Cooling stop temperature: 500°C or less In order to suppress ferrite transformation, rapid cooling is performed to a cooling stop temperature of 500°C or less. The cooling stop temperature is preferably 450°C or less. Although the lower limit of the cooling stop temperature is not particularly limited, it is preferably 100° C. or higher.
上述した焼鈍工程の後,冷延板を伸長率0.5%以上にて圧延して第二冷延板を得る,圧延工程を行う。ここまでの工程で得られた冷延板は可動転位を多く含む。本圧延工程により,可動転位が粒界に集積し絡み合い不動転位となる。伸長率が0.5%未満では効果が小さい。圧延工程における伸長率は,好ましくは0.6%以上とする。圧延工程における伸長率の上限は特に設けないが,設備への負荷をより低減するため,例えば2%以下が好ましい。 After the annealing step, a rolling step in which rolling is performed with an elongation of 0.5% or more After the above-described annealing step, the cold-rolled sheet is rolled at an elongation of 0.5% or more to obtain a second cold-rolled sheet , perform the rolling process. The cold-rolled sheet obtained through the steps up to this point contains many mobile dislocations. In this rolling process, mobile dislocations accumulate at grain boundaries and become entangled to become immobile dislocations. If the elongation rate is less than 0.5%, the effect is small. The elongation rate in the rolling process is preferably 0.6% or more. Although the upper limit of the elongation rate in the rolling process is not particularly set, it is preferably 2% or less, for example, in order to further reduce the load on the equipment.
上述した圧延工程後は,粒界近傍に集積した転位上にCを偏析し,あるいはクラスター生成させるため,低温で第二冷延板を焼戻す。再加熱温度が70℃未満ではCの拡散が遅く,十分な量までCが粒界近傍に濃化しない。一方,再加熱温度が200℃超では過度に焼戻しが進み,セメンタイトが析出する。粒界に析出したセメンタイトは破壊起点になりやすく,またセメンタイト周囲の母相のC濃度が低下するため,曲げ性及び靭性が低下する。再加熱温度は,好ましくは90℃以上とする。また,再加熱温度は,好ましくは190℃以下とする。 After the rolling step, a reheating step in which the second cold-rolled sheet is held at a reheating temperature of 70°C or higher and 200°C or lower for 600 seconds or longer. Alternatively, the second cold-rolled sheet is tempered at a low temperature to form clusters. If the reheating temperature is less than 70° C., diffusion of C is slow and C does not concentrate near the grain boundary to a sufficient amount. On the other hand, if the reheating temperature exceeds 200° C., the tempering proceeds excessively and cementite precipitates. The cementite precipitated at the grain boundary tends to become a fracture initiation point, and the C concentration in the matrix around the cementite decreases, resulting in a decrease in bendability and toughness. The reheating temperature is preferably 90°C or higher. Also, the reheating temperature is preferably 190° C. or lower.
再加熱温度での保持時間が600s未満では,Cの拡散が遅く,十分な量のC濃化が得られない。再加熱温度での保持時間の上限は特に限定しないが,セメンタイトの析出を防ぐために,好ましくは43200s(0.5日)以下である。再加熱温度での保持時間は,好ましくは800s以上である。 Holding time at reheating temperature: 600 s or more If the holding time at reheating temperature is less than 600 s, diffusion of C is slow and a sufficient amount of C concentration cannot be obtained. Although the upper limit of the holding time at the reheating temperature is not particularly limited, it is preferably 43200 seconds (0.5 days) or less in order to prevent precipitation of cementite. The holding time at the reheating temperature is preferably 800 s or more.
上述した高強度鋼板又は高強度めっき鋼板を少なくとも一部に用いてなる部材を提供することができる。上述した高強度鋼板又は高強度めっき鋼板を,一例においてはプレス加工により目的の形状に成形し,自動車部品とすることができる。なお,自動車部品は,本実施形態に係る高強度鋼板または高強度めっき鋼板以外の鋼板を,素材として含んでいてもよい。本実施形態によれば,TSが1180MPa以上であり,曲げ性,靭性,及び高降伏比を兼備した高強度鋼板を提供することができる。そのため,車体の軽量化に寄与する自動車部品として好適である。本高強度鋼板又は高強度めっき鋼板は,自動車部品の中でも,特に,骨格構造部品または補強部品として使用される部材全般において好適に用いることができる。 [Element]
It is possible to provide a member at least partially using the high-strength steel sheet or high-strength plated steel sheet described above. In one example, the high-strength steel sheet or high-strength galvanized steel sheet described above can be formed into a desired shape by press working and used as an automobile part. In addition, the automobile part may contain a steel sheet other than the high-strength steel sheet or the high-strength plated steel sheet according to the present embodiment as a material. According to this embodiment, it is possible to provide a high-strength steel sheet having a TS of 1180 MPa or more and having bendability, toughness and a high yield ratio. Therefore, it is suitable as an automobile part that contributes to weight reduction of the vehicle body. The present high-strength steel sheet or high-strength galvanized steel sheet can be suitably used in general members used as frame structural parts or reinforcing parts, among automobile parts.
得られた鋼板に対しJIS Z 2241に準拠して引張試験を行った。圧延方向と直交方向を長手方向としてJIS5号引張試験片を採取し,引張試験を行って引張強さ(TS)と降伏強度(YS)とを測定した。引張強さTSが1180MPa以上であれば引張強さが良好と判断した。また降伏強度と引張強さとの比YR=YS/TSが0.80以上であれば高降伏比とした。 [Tensile test]
A tensile test was performed on the obtained steel sheets in accordance with JIS Z 2241. A JIS No. 5 tensile test piece was taken with the longitudinal direction perpendicular to the rolling direction, and a tensile test was performed to measure the tensile strength (TS) and yield strength (YS). If the tensile strength TS was 1180 MPa or more, it was judged that the tensile strength was good. Also, if the ratio of yield strength to tensile strength YR=YS/TS is 0.80 or more, it is regarded as a high yield ratio.
シャルピー衝撃試験は,JIS Z 2242に準拠して行った。得られた鋼板より,鋼板の圧延方向に対して直角方向がVノッチ付与方向となるように,幅が10mm,長さが55mm,長さの中央部にノッチ深さが2mmとなるよう90°のVノッチを付与した試験片を採取した。その後,-120~+120℃の試験温度域でシャルピー衝撃試験を行なった。得られた脆性破面率より遷移曲線を求め,脆性破面率が50%となる温度を脆性-延性遷移温度と決定した。なお,シャルピー試験より得られた脆性-延性遷移温度が-40°C以下の場合を靱性が良好と判断した。表中では,脆性-延性遷移温度が-40°C以下の場合を靭性が「優」,脆性-延性遷移温度が-40°C超の場合を靭性が「劣」として示した。 [Charpy test]
The Charpy impact test conforms to JIS Z. 2242. From the obtained steel sheet, a 90° A test piece with a V-notch of was taken. After that, a Charpy impact test was performed in a test temperature range of -120 to +120°C. A transition curve was obtained from the obtained brittle fracture surface ratio, and the temperature at which the brittle fracture surface ratio was 50% was determined as the brittle-ductile transition temperature. In addition, when the brittle-ductile transition temperature obtained from the Charpy test was -40°C or less, the toughness was judged to be good. In the table, when the brittle-ductile transition temperature is −40° C. or less, the toughness is “excellent”, and when the brittle-ductile transition temperature is over −40° C., the toughness is “poor”.
曲げ試験は,JIS Z 2248に準拠して行った。得られた鋼板より,鋼板の圧延方向に対して平行方向が曲げ試験の軸方向となるように,幅が30mm,長さが100mmの短冊状の試験片を採取した。その後,押込み荷重を100kN,押付け保持時間を5秒とする条件で,90°V曲げ試験を行った。なお,曲げ性は曲げ試験の合格率で評価し,曲げ半径(R)を板厚(t)で除した値R/tが5以下となる最大のR(例えば,板厚が1.2mmの場合,曲げ半径は7.0mm)において,5サンプルの曲げ試験を実施し,次いで,曲げ頂点の稜線部における亀裂発生有無の評価を行い,5サンプルとも割れない場合,つまり,合格率100%の場合のみ,曲げ性が良好と判断した。表中では,合格率100%の場合のみ,曲げ性が「優」,その他の場合を曲げ性が「劣」として示した。ここで,亀裂発生有無は,曲げ頂点の稜線部をデジタルマイクロスコープ(RH-2000:株式会社ハイロックス製)を用いて,40倍の倍率で測定することにより評価した。 [Bending test]
The bending test is JIS Z. 2248. A strip-shaped test piece having a width of 30 mm and a length of 100 mm was taken from the obtained steel sheet so that the direction parallel to the rolling direction of the steel sheet was the axial direction of the bending test. After that, a 90° V bending test was performed under the conditions of an indentation load of 100 kN and a pressing holding time of 5 seconds. In addition, bendability is evaluated by the pass rate of the bending test, and the maximum R at which the value R / t obtained by dividing the bending radius (R) by the plate thickness (t) is 5 or less (for example, when the plate thickness is 1.2 mm In the case, the bending radius is 7.0 mm), a bending test is performed on 5 samples, and then the presence or absence of crack generation at the ridge of the bending apex is evaluated. The bendability was judged to be good only in the case. In the table, the bendability is indicated as "excellent" only when the pass rate is 100%, and the bendability is indicated as "poor" in the other cases. Here, the presence or absence of crack generation was evaluated by measuring the ridgeline portion of the bending apex with a digital microscope (RH-2000: manufactured by Hylox Co., Ltd.) at a magnification of 40 times.
Claims (7)
- 質量%で,
C: 0.10%以上0.30以下,
Si: 0.20%以上1.20%以下,
Mn: 2.5%以上4.0%以下,
P: 0.050%以下,
S: 0.020%以下,
Al: 0.10%以下,
N: 0.01%以下,
Ti: 0.100%以下,
Nb: 0.002%以上0.050%以下及び
B: 0.0005%以上0.0050%以下
を含有し,残部がFe及び不可避的不純物からなり,下記式(1)を満足する成分組成を有し,
マルテンサイト及びベイナイトの面積率の合計が95%以上であり,
旧オーステナイト粒の平均結晶粒径が10μm以下であり,
旧オーステナイト粒界のB濃度が質量%で0.10%以上であり,
マルテンサイト粒界に沿ってC濃化領域を有し,
前記C濃化領域のC濃度が鋼中のC含有量の4.0倍以上であり,
前記マルテンサイト粒界と直交する方向において3nm以上100nm以下の濃化幅を有し,かつ前記マルテンサイト粒界に平行な方向において100nm以上の長さを有する,高強度鋼板。
([%N]/14)/([%Ti]/47.9)<1.0…(1)
式(1)において,[%N]及び[%Ti]はそれぞれN及びTiの鋼中含有量(質量%)を示す。 % by mass,
C: 0.10% or more and 0.30 or less,
Si: 0.20% or more and 1.20% or less,
Mn: 2.5% or more and 4.0% or less,
P: 0.050% or less,
S: 0.020% or less,
Al: 0.10% or less,
N: 0.01% or less,
Ti: 0.100% or less,
Nb: 0.002% or more and 0.050% or less and B: 0.0005% or more and 0.0050% or less, with the balance being Fe and unavoidable impurities, having a composition that satisfies the following formula (1): have
The total area ratio of martensite and bainite is 95% or more,
The average crystal grain size of the prior austenite grains is 10 μm or less,
The B concentration of the prior austenite grain boundary is 0.10% or more by mass%,
It has a C-enriched region along the martensite grain boundary,
The C concentration in the C-enriched region is 4.0 times or more the C content in the steel,
A high-strength steel sheet having a condensed width of 3 nm or more and 100 nm or less in a direction perpendicular to the martensite grain boundary and a length of 100 nm or more in a direction parallel to the martensite grain boundary.
([%N]/14)/([%Ti]/47.9)<1.0 (1)
In the formula (1), [%N] and [%Ti] indicate the content (% by mass) of N and Ti in the steel, respectively. - 前記成分組成が,さらに質量%で,
V: 0.100以下,
Mo: 0.500%以下,
Cr: 1.00%以下,
Cu: 1.00%以下,
Ni: 0.50%以下,
Sb: 0.200%以下,
Sn: 0.200%以下,
Ta: 0.200%以下,
W: 0.400%以下,
Zr: 0.0200%以下,
Ca: 0.0200%以下,
Mg: 0.0200%以下,
Co: 0.020%以下,
REM: 0.0200%以下,
Te: 0.020%以下,
Hf: 0.10%以下及び
Bi: 0.200%以下
のうちから選ばれる少なくとも1種の元素を含有する,請求項1に記載の高強度鋼板。 The component composition is further mass%,
V: 0.100 or less,
Mo: 0.500% or less,
Cr: 1.00% or less,
Cu: 1.00% or less,
Ni: 0.50% or less,
Sb: 0.200% or less,
Sn: 0.200% or less,
Ta: 0.200% or less,
W: 0.400% or less,
Zr: 0.0200% or less,
Ca: 0.0200% or less,
Mg: 0.0200% or less,
Co: 0.020% or less,
REM: 0.0200% or less,
Te: 0.020% or less,
The high-strength steel sheet according to claim 1, containing at least one element selected from Hf: 0.10% or less and Bi: 0.200% or less. - 請求項1または2に記載の高強度鋼板の少なくとも片面にめっき層を有する,高強度めっき鋼板。 A high-strength plated steel sheet having a plating layer on at least one side of the high-strength steel sheet according to claim 1 or 2.
- 請求項1または2の成分組成を有する鋼スラブに熱間圧延を施して熱延板とし,
前記熱延板に冷間圧延を施して冷延板とし,
前記冷延板を,850℃以上920℃以下の第一加熱温度まで加熱して10s以上保持し,次いで,1000℃以上1200℃以下の第二加熱温度まで50℃/s以上の平均加熱速度で昇温し,該第二加熱温度に到達後5秒以内に,50℃/s以上の平均冷却速度で500℃以下まで冷却する,焼鈍工程を行い,
前記焼鈍工程の後,前記冷延板を伸長率0.5%以上にて圧延して第二冷延板を得る,圧延工程を行い,
前記圧延工程の後,前記第二冷延板を70℃以上200℃以下の再加熱温度に600s以上保持する再加熱工程を行なって高強度鋼板を得る,高強度鋼板の製造方法。 A steel slab having the composition of claim 1 or 2 is hot-rolled to form a hot-rolled sheet,
cold-rolling the hot-rolled sheet to obtain a cold-rolled sheet;
The cold-rolled sheet is heated to a first heating temperature of 850° C. or higher and 920° C. or lower and held for 10 s or longer, and then heated to a second heating temperature of 1000° C. or higher and 1200° C. or lower at an average heating rate of 50° C./s or higher. An annealing step is performed in which the temperature is raised and cooled to 500 ° C. or less at an average cooling rate of 50 ° C./s or more within 5 seconds after reaching the second heating temperature,
After the annealing step, performing a rolling step of rolling the cold-rolled sheet at an elongation of 0.5% or more to obtain a second cold-rolled sheet,
A method for producing a high-strength steel sheet, wherein after the rolling step, a reheating step is performed in which the second cold-rolled sheet is held at a reheating temperature of 70° C. to 200° C. for 600 seconds or more to obtain a high-strength steel sheet. - 請求項4に記載の焼鈍工程の後,再加熱工程の前に,前記高強度鋼板の少なくとも片面にめっき処理を施して高強度めっき鋼板を得る,めっき工程を有する,高強度めっき鋼板の製造方法。 A method for producing a high-strength plated steel sheet, comprising a plating step of applying a plating treatment to at least one side of the high-strength steel sheet to obtain a high-strength plated steel sheet after the annealing step according to claim 4 and before the reheating step. .
- 請求項1または2に記載の高強度鋼板を少なくとも一部に用いてなる,部材。 A member at least partially using the high-strength steel sheet according to claim 1 or 2.
- 請求項3に記載の高強度めっき鋼板を少なくとも一部に用いてなる,部材。
A member at least partly made of the high-strength plated steel sheet according to claim 3.
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