JP6123957B1 - High strength steel plate and manufacturing method thereof - Google Patents
High strength steel plate and manufacturing method thereof Download PDFInfo
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- JP6123957B1 JP6123957B1 JP2016567261A JP2016567261A JP6123957B1 JP 6123957 B1 JP6123957 B1 JP 6123957B1 JP 2016567261 A JP2016567261 A JP 2016567261A JP 2016567261 A JP2016567261 A JP 2016567261A JP 6123957 B1 JP6123957 B1 JP 6123957B1
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 112
- 239000010959 steel Substances 0.000 title claims abstract description 112
- 238000004519 manufacturing process Methods 0.000 title claims description 27
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 72
- 239000002245 particle Substances 0.000 claims abstract description 29
- 238000001556 precipitation Methods 0.000 claims abstract description 18
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 17
- 239000000203 mixture Substances 0.000 claims abstract description 17
- 229910001562 pearlite Inorganic materials 0.000 claims abstract description 12
- 229910001567 cementite Inorganic materials 0.000 claims abstract description 11
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 9
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 claims abstract description 9
- 238000005096 rolling process Methods 0.000 claims description 40
- 238000000034 method Methods 0.000 claims description 36
- 238000002791 soaking Methods 0.000 claims description 29
- 238000000137 annealing Methods 0.000 claims description 28
- 238000001816 cooling Methods 0.000 claims description 28
- 239000013078 crystal Substances 0.000 claims description 28
- 229910052758 niobium Inorganic materials 0.000 claims description 24
- 229910052719 titanium Inorganic materials 0.000 claims description 24
- 238000005098 hot rolling Methods 0.000 claims description 23
- 229910052720 vanadium Inorganic materials 0.000 claims description 17
- 238000007747 plating Methods 0.000 claims description 16
- 238000005097 cold rolling Methods 0.000 claims description 12
- 238000005275 alloying Methods 0.000 claims description 9
- 230000000717 retained effect Effects 0.000 claims description 9
- 239000010960 cold rolled steel Substances 0.000 claims description 8
- 238000004804 winding Methods 0.000 claims description 8
- 229910052799 carbon Inorganic materials 0.000 claims description 7
- 238000005246 galvanizing Methods 0.000 claims description 7
- 229910052757 nitrogen Inorganic materials 0.000 claims description 7
- 239000012535 impurity Substances 0.000 claims description 6
- 229910052804 chromium Inorganic materials 0.000 claims description 5
- 229910052802 copper Inorganic materials 0.000 claims description 5
- 229910052748 manganese Inorganic materials 0.000 claims description 4
- 229910052750 molybdenum Inorganic materials 0.000 claims description 3
- 229910052698 phosphorus Inorganic materials 0.000 claims description 3
- 229910052759 nickel Inorganic materials 0.000 claims description 2
- 238000009864 tensile test Methods 0.000 description 12
- 230000000694 effects Effects 0.000 description 10
- 150000001247 metal acetylides Chemical class 0.000 description 10
- 239000002244 precipitate Substances 0.000 description 8
- 230000007423 decrease Effects 0.000 description 7
- 238000005266 casting Methods 0.000 description 6
- 239000002131 composite material Substances 0.000 description 6
- 239000011573 trace mineral Substances 0.000 description 6
- 235000013619 trace mineral Nutrition 0.000 description 6
- 238000010438 heat treatment Methods 0.000 description 5
- 238000001953 recrystallisation Methods 0.000 description 5
- 239000006104 solid solution Substances 0.000 description 5
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 3
- 239000000463 material Substances 0.000 description 3
- 238000005204 segregation Methods 0.000 description 3
- 238000005728 strengthening Methods 0.000 description 3
- 229910052787 antimony Inorganic materials 0.000 description 2
- 239000011248 coating agent Substances 0.000 description 2
- 238000000576 coating method Methods 0.000 description 2
- 230000007797 corrosion Effects 0.000 description 2
- 238000005260 corrosion Methods 0.000 description 2
- 238000000605 extraction Methods 0.000 description 2
- 230000003287 optical effect Effects 0.000 description 2
- 238000007670 refining Methods 0.000 description 2
- 230000009466 transformation Effects 0.000 description 2
- 230000032683 aging Effects 0.000 description 1
- 229910001563 bainite Inorganic materials 0.000 description 1
- 230000015572 biosynthetic process Effects 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 238000007796 conventional method Methods 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
- 238000000151 deposition Methods 0.000 description 1
- 230000006866 deterioration Effects 0.000 description 1
- 230000002349 favourable effect Effects 0.000 description 1
- 238000005244 galvannealing Methods 0.000 description 1
- 238000000227 grinding Methods 0.000 description 1
- 230000001771 impaired effect Effects 0.000 description 1
- 230000014759 maintenance of location Effects 0.000 description 1
- 239000000155 melt Substances 0.000 description 1
- 239000002184 metal Substances 0.000 description 1
- 229910052751 metal Inorganic materials 0.000 description 1
- 238000005554 pickling Methods 0.000 description 1
- 238000003825 pressing Methods 0.000 description 1
- 238000012545 processing Methods 0.000 description 1
- 238000011160 research Methods 0.000 description 1
- 238000005482 strain hardening Methods 0.000 description 1
- 239000000126 substance Substances 0.000 description 1
- 238000004381 surface treatment Methods 0.000 description 1
- 238000012360 testing method Methods 0.000 description 1
- 239000010409 thin film Substances 0.000 description 1
- 229910052718 tin Inorganic materials 0.000 description 1
- 229910052721 tungsten Inorganic materials 0.000 description 1
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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
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
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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
- C21D11/00—Process control or regulation for heat treatments
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- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
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- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0236—Cold rolling
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- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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- C—CHEMISTRY; METALLURGY
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- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
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- 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
- C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
- C21D9/54—Furnaces for treating strips or wire
- C21D9/56—Continuous furnaces for strip or wire
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
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- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
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- 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
- C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
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- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
- C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
- C23C2/0224—Two or more thermal pretreatments
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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
- C23C2/06—Zinc or cadmium or alloys based thereon
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- 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
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- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
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- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
- C23C2/36—Elongated material
- C23C2/40—Plates; Strips
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- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
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Abstract
高降伏比で引張特性の異方性が小さい高強度鋼板を得る。特定の成分組成とし、鋼組織は、面積率でフェライト:90%以上、パーライトとセメンタイトの合計:0〜10%、マルテンサイトと残留オーステナイトの合計:0〜2%からなり、前記フェライトの平均結晶粒径が15.0μm以下、平均アスペクト比が5.0以下であり、Nb炭化物、Ti炭化物および/またはV炭化物を含み、該Nb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径が5〜50nmであり、Nb炭化物、Ti炭化物およびV炭化物の析出量の合計が体積率で0.005〜0.050%である高強度鋼板とする。A high strength steel sheet having a high yield ratio and small anisotropy in tensile properties is obtained. The steel structure has a specific component composition, and the steel structure is composed of ferrite in an area ratio of 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and residual austenite: 0 to 2%. The particle size is 15.0 μm or less, the average aspect ratio is 5.0 or less, and contains Nb carbide, Ti carbide and / or V carbide, and the average particle size of the Nb carbide and Ti carbide and / or V carbide is 5 to 5 The high-strength steel sheet is 50 nm, and the total precipitation amount of Nb carbide, Ti carbide, and V carbide is 0.005 to 0.050% by volume.
Description
本発明は、自動車部品等に適用される高強度鋼板およびその製造方法に関する。 The present invention relates to a high-strength steel plate applied to automobile parts and the like and a method for manufacturing the same.
自動車部品等の素材として、素材の薄肉化による部品軽量化などの観点から、高強度鋼板が好ましく用いられる。例えば、骨格用部品や耐衝突用部品などでは、乗員の安全確保のため衝突時に変形しにくいこと、すなわち高い降伏比が要求される。一方、割れが発生することなく安定的にプレス成形するために、引張特性が鋼板内で均一であること、すなわち引張特性の異方性が小さい高強度鋼板が望まれる。このような要求に対し、これまでに種々の鋼板およびその製造技術が開示されている。 A high strength steel plate is preferably used as a material for automobile parts and the like from the viewpoint of reducing the weight of the component by reducing the thickness of the material. For example, skeletal parts and anti-collision parts are required to be difficult to be deformed at the time of collision, that is, to have a high yield ratio in order to ensure the safety of passengers. On the other hand, in order to stably press-form without generating cracks, a high-strength steel sheet having a uniform tensile property within the steel sheet, that is, a small anisotropy of the tensile property is desired. In response to such demands, various steel plates and manufacturing techniques thereof have been disclosed so far.
特許文献1には、Nb、Tiを合計で0.01質量%以上含有し、再結晶率80%以上のフェライトを主相とする高強度鋼板とその製造方法が開示されている。 Patent Document 1 discloses a high-strength steel sheet containing Nb and Ti in a total amount of 0.01% by mass or more and containing ferrite having a recrystallization rate of 80% or more as a main phase and a method for manufacturing the same.
また、特許文献2には、鋼組織に20〜50面積%の未再結晶フェライトを含む耐衝突特性に優れた高強度鋼板とその製造方法が開示されている。 Patent Document 2 discloses a high-strength steel sheet excellent in impact resistance and containing 20-50 area% of non-recrystallized ferrite in the steel structure and a method for producing the same.
特許文献3には、V、Ti、Nbの1種または2種以上を添加し、主相がフェライトまたはベイナイトで粒界における鉄炭化物の析出量を一定以下に制限し、かつ該鉄炭化物の最大粒子径を1μm以下に制御する溶融めっき高強度鋼板とその製造方法が開示されている。 In Patent Document 3, one or more of V, Ti, and Nb is added, the main phase is ferrite or bainite, the amount of precipitation of iron carbide at the grain boundary is limited to a certain value, and the maximum of the iron carbide A hot-dip hot-dip steel sheet that controls the particle diameter to 1 μm or less and a manufacturing method thereof are disclosed.
しかしながら、特許文献1の技術では、熱間圧延後〜650℃の保持時間と連続焼鈍炉での均熱後の冷却における500〜400℃の保持時間のいずれも制御しておらず、本発明で重要なNb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径を制御できていないと考えられるので、高降伏比で引張特性の異方性が小さい高強度鋼板を得ることができない。 However, in the technique of Patent Document 1, neither the holding time of 650 ° C. after hot rolling nor the holding time of 500-400 ° C. in cooling after soaking in a continuous annealing furnace is controlled. Since it is considered that the average particle diameter of important Nb carbide and Ti carbide and / or V carbide cannot be controlled, a high-strength steel sheet having a high yield ratio and small anisotropy in tensile properties cannot be obtained.
特許文献2の技術では、NbやTiを多量添加し、鋼組織として未再結晶フェライトを面積率で20%以上含むため、所望の降伏強さと引張強さの範囲を外れるばかりか、未再結晶フェライトが鋼板組織中に不均一に分散するため局所的に降伏しやすくなり、高降伏比で引張特性の異方性が小さい高強度鋼板が得られない。 In the technique of Patent Document 2, a large amount of Nb or Ti is added, and the steel structure contains unrecrystallized ferrite in an area ratio of 20% or more, so that not only the desired yield strength and tensile strength are out of the range, but also unrecrystallized. Since ferrite is unevenly dispersed in the steel sheet structure, it is easy to yield locally, and a high-strength steel sheet with a high yield ratio and small anisotropy in tensile properties cannot be obtained.
特許文献3の技術では、熱間圧延後〜650℃の保持時間と連続焼鈍炉での均熱後の冷却における500〜400℃の保持時間のいずれも制御しておらず、本発明で重要なNb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径を制御できていないと考えられるので、高降伏比であるとともに引張特性の異方性が小さい高強度鋼板は得らない。 In the technique of Patent Document 3, neither the holding time of 650 ° C. after hot rolling nor the holding time of 500-400 ° C. in cooling after soaking in a continuous annealing furnace is controlled, which is important in the present invention. Since it is considered that the average particle diameter of Nb carbide and Ti carbide and / or V carbide cannot be controlled, a high-strength steel sheet having a high yield ratio and small anisotropy in tensile properties cannot be obtained.
本発明は、このような事情を鑑みてなされたものであり、その目的は、高降伏比であるとともに引張特性の異方性が小さい高強度鋼板を得ることである。 The present invention has been made in view of such circumstances, and an object thereof is to obtain a high-strength steel sheet having a high yield ratio and a small anisotropy of tensile properties.
本発明者らは、上記の課題を解決すべく鋭意研究を実施した。その結果、フェライトを主体とする鋼組織において、フェライトの平均結晶粒径を一定以下に微細化し、平均アスペクト比が一定以下である等軸な組織とし、さらに、Nb炭化物並びにTi炭化物および/またはV炭化物の体積率と粒子径を適正に制御することが重要であることを見出した。そして、所望の組織等に調整するためには、所定の成分組成に調整するとともに、熱間圧延後の巻取温度、焼鈍の昇温時所定温度域での滞留時間と均熱温度を適正な範囲に制御することが有効であることを見出した。 The inventors of the present invention conducted intensive research to solve the above problems. As a result, in the steel structure mainly composed of ferrite, the average crystal grain size of ferrite is refined to a certain level or less, an equiaxed structure having an average aspect ratio of a certain level or less, Nb carbide and / or Ti carbide and / or V It was found that it is important to properly control the volume fraction and particle size of the carbide. And in order to adjust to a desired structure etc., while adjusting to a predetermined component composition, the coiling temperature after hot rolling, the residence time and the soaking temperature in the predetermined temperature range at the time of temperature rise of annealing are appropriate. We found that it was effective to control the range.
本発明は、以上の知見に基づきなされたもので、その要旨は以下のとおりである。 The present invention has been made based on the above findings, and the gist thereof is as follows.
[1]成分組成は、質量%で、C:0.02%〜0.10%未満、Si:0.10%未満、Mn:1.0%未満、P:0.10%以下、S:0.020%以下、Al:0.01〜0.10%、N:0.010%以下、Nb:0.005〜0.070%、Ti:0.100%以下(0%を含む)、V:0.100%以下(0%を含む)かつNbとTiとVを合計で0.005〜0.100%を含有し、残部がFeおよび不可避的不純物からなり、鋼組織は、面積率でフェライト:90%以上、パーライトとセメンタイトの合計:0〜10%、マルテンサイトと残留オーステナイトの合計:0〜3%からなり、前記フェライトの平均結晶粒径が15.0μm以下、平均アスペクト比が5.0以下であり、Nb炭化物並びにTi炭化物および/またはV炭化物を含み、該Nb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径が5〜50nmであり、Nb炭化物、Ti炭化物およびV炭化物の析出量の合計が体積率で0.005〜0.050%である高強度鋼板。 [1] Component composition is mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.10% or less, S: 0.020% or less, Al: 0.01 to 0.10%, N: 0.010% or less, Nb: 0.005 to 0.070%, Ti: 0.100% or less (including 0%), V: 0.100% or less (including 0%) and Nb, Ti and V in total 0.005 to 0.100%, with the balance being Fe and inevitable impurities, and the steel structure has an area ratio Ferrite: 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and residual austenite: 0 to 3%, average grain size of the ferrite is 15.0 μm or less, average aspect ratio is 5.0 or less, Nb carbide and Ti carbide and / or The Nb carbide and Ti carbide and / or V carbide have an average particle diameter of 5 to 50 nm, and the total amount of precipitation of Nb carbide, Ti carbide and V carbide is 0.005 to 0.050 in volume ratio. % High-strength steel sheet.
[2]前記成分組成は、さらに、質量%で、Cr:0.3%以下、Mo:0.3%以下、B:0.005%以下、Cu:0.3%以下、Ni:0.3%以下、Sb:0.3%以下のいずれか1種または2種以上を含有する[1]に記載の高強度鋼板。 [2] The component composition further includes, by mass%, Cr: 0.3% or less, Mo: 0.3% or less, B: 0.005% or less, Cu: 0.3% or less, Ni: 0.00. The high-strength steel sheet according to [1], containing any one or more of 3% or less and Sb: 0.3% or less.
[3]表面に亜鉛めっき層を有する[1]または[2]に記載の高強度鋼板。 [3] The high-strength steel plate according to [1] or [2], which has a galvanized layer on the surface.
[4]前記亜鉛めっき層が溶融亜鉛めっき層である[3]に記載の高強度鋼板。 [4] The high-strength steel plate according to [3], wherein the galvanized layer is a hot-dip galvanized layer.
[5]前記溶融亜鉛めっき層が合金化溶融亜鉛めっき層である[4]に記載の高強度鋼板。 [5] The high-strength steel sheet according to [4], wherein the hot-dip galvanized layer is an alloyed hot-dip galvanized layer.
[6]前記亜鉛めっき層が電気亜鉛めっき層である[3]に記載の高強度鋼板。 [6] The high-strength steel plate according to [3], wherein the galvanized layer is an electrogalvanized layer.
[7][1]または[2]に記載の高強度鋼板の製造方法であって、鋼を熱間圧延し、該熱間圧延後、仕上圧延温度〜650℃の温度域の滞留時間を10秒以下の条件で鋼板を冷却し、500〜700℃で巻取る熱間圧延工程と、前記熱間圧延工程で得られる熱延鋼板を75%以下の圧延率で冷間圧延する冷間圧延工程と、前記冷間圧延工程で得られる冷延鋼板を、連続焼鈍炉で、昇温時における650〜750℃の温度域で滞留時間:60秒以下で滞留し、該滞留後に均熱温度:760〜880℃、均熱時間:120秒以下の条件で均熱し、400〜500℃の温度域の滞留時間が100秒以下の条件で冷却する焼鈍工程と、を有する高強度鋼板の製造方法。 [7] A method for producing a high-strength steel sheet according to [1] or [2], wherein the steel is hot-rolled, and after the hot-rolling, a residence time in a temperature range of finish rolling temperature to 650 ° C. is 10 A hot rolling process in which the steel sheet is cooled under conditions of less than a second and wound at 500 to 700 ° C., and a cold rolling process in which the hot rolled steel sheet obtained in the hot rolling process is cold rolled at a rolling rate of 75% or less. Then, the cold-rolled steel sheet obtained in the cold rolling step is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at the time of temperature rise in a residence time: 60 seconds or less, and after the residence, a soaking temperature: 760 -880 degreeC, soaking time: The manufacturing method of the high strength steel plate which has soaking | uniform-heating on the conditions for 120 second or less, and cooling on the conditions whose residence time of 400-500 degreeC is 100 seconds or less.
[8]前記焼鈍工程後の冷延鋼板を、めっき処理するめっき工程を有する[7]に記載の高強度鋼板の製造方法。 [8] The method for producing a high-strength steel plate according to [7], including a plating step of plating the cold-rolled steel plate after the annealing step.
[9]前記めっき処理は、溶融亜鉛めっき処理である[8]に記載の高強度鋼板の製造方法。 [9] The method for manufacturing a high-strength steel sheet according to [8], wherein the plating treatment is a hot dip galvanizing treatment.
[10]前記めっき工程後の冷延鋼板を、合金化処理する合金化工程を有する[9]に記載の高強度鋼板の製造方法。 [10] The method for producing a high-strength steel plate according to [9], including an alloying step of alloying the cold-rolled steel plate after the plating step.
[11]前記めっき処理は、電気亜鉛めっき処理である[8]に記載の高強度鋼板の製造方法。 [11] The method for producing a high-strength steel sheet according to [8], wherein the plating treatment is an electrogalvanizing treatment.
本発明では、成分組成、熱延後の巻取条件、焼鈍の昇温時所定温度域での滞留時間と均熱温度などの製造条件を適正に制御する。この制御により、本発明が目的とする鋼組織が得られ、その結果、自動車部品等の用途に要求される高降伏比で引張特性の異方性が小さい高強度鋼板を安定して製造することが可能になる。本発明の高強度鋼板により、自動車の更なる軽量化が可能になり、本発明の自動車、鉄鋼業界における利用価値は極めて大きい。 In the present invention, the production conditions such as the component composition, the winding condition after hot rolling, the residence time in the predetermined temperature range at the time of temperature rise of annealing, and the soaking temperature are appropriately controlled. By this control, the steel structure intended by the present invention can be obtained, and as a result, a high strength steel sheet having a high yield ratio and a small anisotropy in tensile properties required for applications such as automobile parts can be stably produced. Is possible. The high-strength steel sheet of the present invention makes it possible to further reduce the weight of an automobile, and the utility value of the present invention in the automobile and steel industry is extremely large.
以下、本発明の実施形態について説明する。なお、本発明は以下の実施形態に限定されない。 Hereinafter, embodiments of the present invention will be described. In addition, this invention is not limited to the following embodiment.
先ず、本発明の高強度鋼板の概要について説明する。 First, the outline | summary of the high strength steel plate of this invention is demonstrated.
本発明の高強度鋼板は、330MPa〜500MPa未満の引張強さ、0.70以上の降伏比を有し、引張方向が圧延方向と平行な向きで行った引張試験における降伏強さと、引張方向が圧延方向と垂直な向きで行った引張試験における降伏強さの差が30MPa以下である。降伏比が0.70以上であることから、本発明の高強度鋼板は高い降伏比を有する。また、上記降伏強さの差が30MPa以下であることから、本発明の高強度鋼板は、引張特性の異方性が小さい。 The high-strength steel sheet of the present invention has a tensile strength of 330 MPa to less than 500 MPa, a yield ratio of 0.70 or more, and a tensile strength in a tensile test performed in a direction in which the tensile direction is parallel to the rolling direction. The difference in yield strength in a tensile test conducted in a direction perpendicular to the rolling direction is 30 MPa or less. Since the yield ratio is 0.70 or more, the high-strength steel sheet of the present invention has a high yield ratio. Moreover, since the difference in yield strength is 30 MPa or less, the high-strength steel sheet of the present invention has a small anisotropy in tensile properties.
本発明では、Nb:0.005%〜0.070%、Ti:0.100%以下(0%を含む)、V:0.100%以下(0%を含む)、かつNbとTiとVを合計で0.005%〜0.100%以下含有する成分組成とすることが特に重要である。 In the present invention, Nb: 0.005% to 0.070%, Ti: 0.100% or less (including 0%), V: 0.100% or less (including 0%), and Nb, Ti, and V It is particularly important to obtain a component composition that contains 0.005% to 0.100% or less in total.
成分組成や製造条件の調整により、鋼組織を、必須のフェライトと任意のパーライト等で構成し、該フェライトの平均結晶粒径が15.0μm以下、平均アスペクト比が5.0以下、Nb炭化物並びにTi炭化物および/またはV炭化物を含み、該Nb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径が5〜50nmであり、Nb炭化物、Ti炭化物およびV炭化物の析出量の合計が体積率で0.005〜0.050%になるように調整することで、高降伏比で引張特性の異方性が小さい高強度鋼板を得られる。 By adjusting the component composition and manufacturing conditions, the steel structure is composed of essential ferrite and optional pearlite, and the ferrite has an average crystal grain size of 15.0 μm or less, an average aspect ratio of 5.0 or less, Nb carbide, Ti carbide and / or V carbide is included, the average particle diameter of the Nb carbide and Ti carbide and / or V carbide is 5 to 50 nm, and the total precipitation amount of Nb carbide, Ti carbide and V carbide is 0 by volume By adjusting to 0.005 to 0.050%, a high-strength steel sheet having a high yield ratio and small anisotropy in tensile properties can be obtained.
本発明においてNb炭化物、Ti炭化物およびV炭化物にはNb炭窒化物、Ti炭窒化物、V炭窒化物およびNb、Ti複合炭窒化物、Nb、V複合炭窒化物およびNb、Ti、V複合炭窒化物も含む。なお、Nb、Ti複合炭窒化物は、Nbの炭化物と捉えるかTiの炭化物と捉えて平均粒子径や合計体積率を考えればよい。Nb、V複合炭窒化物およびNb、Ti、V複合炭窒化物についても同様である。 In the present invention, Nb carbide, Ti carbide and V carbide include Nb carbonitride, Ti carbonitride, V carbonitride and Nb, Ti composite carbonitride, Nb, V composite carbonitride and Nb, Ti, V composite. Includes carbonitrides. The Nb and Ti composite carbonitrides may be regarded as Nb carbides or Ti carbides, and the average particle diameter and the total volume ratio may be considered. The same applies to Nb, V composite carbonitride and Nb, Ti, V composite carbonitride.
上記の通り、フェライトの平均結晶粒径、平均アスペクト比、および炭化物(Nb炭化物並びにTi炭化物および/またはV炭化物)の平均粒子径と析出量が所望の条件を満たすためには、成分組成のみならず製造条件も重要である。具体的には、熱間圧延後の冷却において仕上圧延温度〜650℃の温度域の滞留時間を10秒以下とし、巻取温度を500〜700℃とする。また、焼鈍の加熱において650〜750℃の温度域での滞留時間を60秒以下とし、引き続き760〜880℃の均熱温度で120秒以下均熱する。巻取後の冷却中にNb炭化物、Ti炭化物および/またはV炭化物を均一微細に析出させ、冷間圧延後、焼鈍にてフェライトを比較的低温で再結晶させ、かつ再結晶温度域での滞留時間を所定以下に制御することで本発明が目的とする鋼組織が得られる。 As described above, in order for the average grain size of ferrite, the average aspect ratio, and the average particle size and precipitation amount of carbides (Nb carbide and Ti carbide and / or V carbide) to satisfy the desired conditions, only the component composition is required. The manufacturing conditions are also important. Specifically, in the cooling after hot rolling, the residence time in the temperature range of the finish rolling temperature to 650 ° C. is set to 10 seconds or less, and the winding temperature is set to 500 to 700 ° C. Further, in the annealing heating, the residence time in the temperature range of 650 to 750 ° C. is set to 60 seconds or less, and then the temperature is soaked at 760 to 880 ° C. for 120 seconds or less. During cooling after winding, Nb carbide, Ti carbide and / or V carbide are uniformly and finely precipitated, and after cold rolling, ferrite is recrystallized at a relatively low temperature and stays in the recrystallization temperature range. By controlling the time below a predetermined value, the steel structure intended by the present invention can be obtained.
降伏強さと引張強さは、引張方向が圧延方向と垂直になるようJIS5号引張試験片を採取し、JIS Z 2241に準拠した引張試験により求める。引張特性の異方性は、引張方向が圧延方向と垂直および平行になる向きからそれぞれJIS5号引張試験片を採取し、JIS Z 2241に準拠した引張試験を行い、降伏強さの差より求める。 The yield strength and the tensile strength are obtained by taking a JIS No. 5 tensile test piece so that the tensile direction is perpendicular to the rolling direction and performing a tensile test in accordance with JIS Z 2241. The anisotropy of the tensile properties is obtained from the difference in yield strength by collecting JIS No. 5 tensile test pieces from the direction in which the tensile direction is perpendicular and parallel to the rolling direction and performing a tensile test in accordance with JIS Z 2241.
以上の知見に基づき完成された本発明の高強度鋼板は、自動車部品等の素材に求められる、高降伏比で引張特性の異方性が小さいという特徴を有する。 The high-strength steel sheet of the present invention completed on the basis of the above knowledge has the characteristics that the anisotropy of tensile properties is small at a high yield ratio required for materials such as automobile parts.
次に、本発明の成分組成の限定理由、鋼組織の限定理由および製造条件の限定理由について説明する。 Next, the reason for limiting the composition of the present invention, the reason for limiting the steel structure, and the reason for limiting the manufacturing conditions will be described.
(1)成分組成
本発明の高強度鋼板は、質量%で、C:0.02%〜0.10%未満、Si:0.10%未満、Mn:1.0%未満、P:0.10%以下、S:0.020%以下、Al:0.01〜0.10%、N:0.010%以下、Nb:0.005〜0.070%、Ti:0.100%以下(0%を含む)およびV:0.100%以下(0%を含む)かつNbとTiとVを合計で0.005〜0.100%以下含有する。(1) Component composition The high-strength steel sheet of the present invention is in mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.00. 10% or less, S: 0.020% or less, Al: 0.01 to 0.10%, N: 0.010% or less, Nb: 0.005 to 0.070%, Ti: 0.100% or less ( 0: inclusive) and V: 0.100% or less (including 0%), and Nb, Ti, and V are contained in a total amount of 0.005 to 0.100% or less.
また、本発明の高強度鋼板は、任意成分として、さらに、質量%で、Cr:0.3%以下、Mo:0.3%以下、B:0.005%以下、Cu:0.3%以下、Ni:0.3%以下、Sb:0.3%以下のいずれか1種または2種以上を含有してもよい。 Moreover, the high-strength steel sheet of the present invention may further include, as an optional component, in mass%, Cr: 0.3% or less, Mo: 0.3% or less, B: 0.005% or less, Cu: 0.3% Hereinafter, any one or more of Ni: 0.3% or less and Sb: 0.3% or less may be contained.
上記以外の残部はFe及び不可避的不純物である。 The balance other than the above is Fe and inevitable impurities.
以下の成分組成の説明において「%」は「質量%」を意味する。 In the following description of the component composition, “%” means “% by mass”.
C:0.02%〜0.10%未満
Cは、Nb炭化物やTi炭化物やV炭化物となったり、パーライトやマルテンサイトを増加させたりすることから、降伏強さと引張強さの増加に有効な元素である。C含有量が0.02%未満では、炭化物の合計析出量が所望の範囲にならないので本発明が目的とする引張強さが得られない。C含有量が0.10%以上になると、パーライト又はマルテンサイトが過度に生成するために降伏比が低下し、引張特性の異方性が増大する。このため、C含有量は0.02%〜0.10%未満とする。好ましくは0.02〜0.06%である。C: 0.02% to less than 0.10% C is effective for increasing yield strength and tensile strength because it becomes Nb carbide, Ti carbide, V carbide, or increases pearlite and martensite. It is an element. If the C content is less than 0.02%, the total precipitation amount of carbides does not fall within the desired range, so that the intended tensile strength of the present invention cannot be obtained. When the C content is 0.10% or more, pearlite or martensite is excessively generated, so that the yield ratio is lowered and the anisotropy of tensile properties is increased. For this reason, C content shall be 0.02%-less than 0.10%. Preferably it is 0.02 to 0.06%.
Si:0.10%未満
Siは、一般にフェライトの固溶強化により降伏強さと引張強さを増加させるのに有効である。しかし、Siを添加すると、加工硬化能の顕著な向上により降伏強さに比べて引張強さの増加量が大きくなり、降伏比が低下し、表面性状が劣化する。このため、Si含有量は0.10%未満とする。なお、Si含有量の下限は特に限定されないが、降伏強さや引張強さはSi以外の構成でも高められるため、本発明では、Si含有量は少ないほど好ましい。したがって、本発明ではSiを添加しなくてもよいが、製造上Siを不可避的に0.005%含む場合がある。Si: Less than 0.10% Si is generally effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. However, when Si is added, the increase in tensile strength is greater than the yield strength due to a significant improvement in work hardening ability, the yield ratio is lowered, and the surface properties are degraded. For this reason, Si content shall be less than 0.10%. In addition, although the lower limit of Si content is not specifically limited, Since yield strength and tensile strength are raised also in structures other than Si, in this invention, it is so preferable that Si content is small. Therefore, in the present invention, Si may not be added, but Si may inevitably be contained in an amount of 0.005% in production.
Mn:1.0%未満
Mnは、フェライトの固溶強化により降伏強さと引張強さを増加させるのに有効である。しかし、Mn含有量が1.0%以上になると、鋼組織中のマルテンサイト分率が増加するため引張強さが過度に増大し、本発明が目的とする引張強さが得られず、降伏比が低下する。このためMn含有量は1.0%未満とする。Mnは添加しなくても良いが、Mnを添加する場合には、下限について好ましいMn含有量は0.2%以上である。上限について好ましいMn含有量は0.8%以下である。Mn: less than 1.0% Mn is effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. However, if the Mn content is 1.0% or more, the martensite fraction in the steel structure increases, so the tensile strength increases excessively, and the intended tensile strength of the present invention cannot be obtained, yielding. The ratio decreases. Therefore, the Mn content is less than 1.0%. Mn may not be added, but when Mn is added, the preferable Mn content for the lower limit is 0.2% or more. A preferable Mn content for the upper limit is 0.8% or less.
P:0.10%以下
Pはフェライトの固溶強化により降伏強さと引張強さを増加させるのに有効である。このため、本発明ではPを適宜含有することができる。しかし、P含有量が0.10%を超えると、鋳造偏析やフェライト粒界偏析によりフェライトの降伏が局所的に起こるようになるため引張特性の異方性が増大する。このためP含有量は0.10%以下とする。Pは添加しなくても良いが、Pを添加する場合には、下限について好ましいP含有量は0.01%以上である。上限について好ましいP含有量は0.04%以下である。P: 0.10% or less P is effective in increasing yield strength and tensile strength by solid solution strengthening of ferrite. For this reason, in this invention, P can be contained suitably. However, if the P content exceeds 0.10%, ferrite yields locally due to casting segregation or ferrite grain boundary segregation, so the anisotropy of tensile properties increases. Therefore, the P content is 0.10% or less. P may not be added, but when P is added, the preferable P content for the lower limit is 0.01% or more. A preferable P content for the upper limit is 0.04% or less.
S:0.020%以下
Sは不純物として不可避的に含まれる元素である。MnSなどの介在物の形成により曲げ性や局部伸びなどの成形性が低下するので、S含有量はできるだけ低減することが好ましい。本発明においてS含有量は0.020%以下とする。好ましくは0.015%以下とする。なお、上記の通りS含有量は低いほど好ましく、本発明ではSを添加しなくてもよい。しかし、製造上Sを0.0003%含む場合がある。S: 0.020% or less S is an element inevitably included as an impurity. Since the formability such as bendability and local elongation is lowered by the formation of inclusions such as MnS, the S content is preferably reduced as much as possible. In the present invention, the S content is 0.020% or less. Preferably, the content is 0.015% or less. In addition, as above-mentioned, S content is so preferable that it is low, and it is not necessary to add S in this invention. However, there are cases in which S is contained in an amount of 0.0003%.
Al:0.01〜0.10%
Alは精錬工程での脱酸のため、また、固溶NをAlNとして固定させるために添加される。十分な効果を得るにはAl含有量を0.01%以上にする必要がある。また、Al含有量が0.10%を超えるとAlNが多量に析出してフェライトの平均アスペクト比が増大して引張特性の異方性の増大を招く。したがってAl含有量は0.01〜0.10%とする。好ましくは0.01〜0.07%とする。また、さらに好ましくは0.01〜0.06%とする。Al: 0.01-0.10%
Al is added for deoxidation in the refining process and for fixing the solid solution N as AlN. In order to obtain a sufficient effect, the Al content needs to be 0.01% or more. On the other hand, if the Al content exceeds 0.10%, a large amount of AlN precipitates and the average aspect ratio of ferrite increases, leading to an increase in anisotropy of tensile properties. Therefore, the Al content is set to 0.01 to 0.10%. Preferably, the content is 0.01 to 0.07%. More preferably, the content is 0.01 to 0.06%.
N:0.010%以下
Nは溶銑の精錬工程までに不可避的に混入する元素である。N含有量が0.010%を超えると、鋳造時にNb炭化物やTi炭化物やV炭化物が析出後、スラブ加熱時にNb炭化物やTi炭化物やV炭化物が溶解せず粗大な炭化物として残留するためフェライト平均結晶粒の粗大化を招く。よってN含有量は0.010%以下とする。なお、本発明ではNを添加しなくてもよいが、製造上Nを0.0005%含む場合がある。N: 0.010% or less N is an element inevitably mixed up to the hot metal refining process. If the N content exceeds 0.010%, Nb carbide, Ti carbide and V carbide will precipitate during casting, and Nb carbide, Ti carbide and V carbide will not dissolve and remain as coarse carbide during slab heating. Causes coarsening of crystal grains. Therefore, the N content is 0.010% or less. In the present invention, N may not be added, but may contain 0.0005% N in production.
Nb:0.005〜0.070%
Nbはフェライト平均結晶粒の微細化、Nb炭化物の析出による降伏比の増加に寄与する重要な元素である。Nb含有量が0.005%未満では、炭化物の析出量が不十分となる場合がある等により、本発明の効果が得られない。また、Nb含有量が0.070%を超えるとNb炭化物が過剰に析出して焼鈍後も延性に乏しい未再結晶フェライトが残存するか、フェライトの平均アスペクト比が5.0を超えるため引張特性の異方性が増大する。したがって、Nb含有量は0.005〜0.070%とする。好ましくは0.005〜0.040%である。Nb: 0.005 to 0.070%
Nb is an important element that contributes to refinement of ferrite average crystal grains and an increase in yield ratio due to precipitation of Nb carbides. If the Nb content is less than 0.005%, the effect of the present invention cannot be obtained due to the fact that the amount of precipitated carbide may be insufficient. Further, if the Nb content exceeds 0.070%, Nb carbides are excessively precipitated and unrecrystallized ferrite having poor ductility remains after annealing, or the average aspect ratio of ferrite exceeds 5.0, so that tensile properties are exceeded. Increases the anisotropy. Therefore, the Nb content is 0.005 to 0.070%. Preferably it is 0.005 to 0.040%.
Ti:0.100%以下(0%を含む)
V:0.100%以下(0%を含む)
NbとTiとVを合計で0.005〜0.100%
TiおよびVは、Nbと複合添加することにより、Ti炭化物および/またはV炭化物として析出し、フェライトの平均アスペクト比を5.0以下に制御するのに寄与する。NbとTiとVの合計が0.005%未満では、Nb炭化物並びにTi炭化物および/またはV炭化物の体積率が不十分となる結果、炭化物の析出量が所望の範囲にならず、本発明の効果が得られない。また、NbとTiとVの合計が0.100%超ではTi炭化物および/またはV炭化物が過剰に析出して焼鈍後も延性に乏しい未再結晶フェライトが残存するため引張特性の異方性が増大する。したがって、TiおよびVは、Ti:0.100%以下(0%を含む)およびV:0.100%以下(0%を含む)かつNbとTiとVを合計で0.005〜0.100%とする。好ましい合計量は0.007〜0.040%とする。Ti: 0.100% or less (including 0%)
V: 0.100% or less (including 0%)
Nb, Ti and V in total 0.005 to 0.100%
When Ti and V are added together with Nb, Ti and V precipitate as Ti carbide and / or V carbide, and contribute to controlling the average aspect ratio of ferrite to 5.0 or less. If the total of Nb, Ti, and V is less than 0.005%, the volume ratio of Nb carbide and Ti carbide and / or V carbide becomes insufficient. As a result, the amount of precipitation of carbide does not fall within the desired range, and The effect is not obtained. Further, if the total of Nb, Ti, and V exceeds 0.100%, Ti carbide and / or V carbide precipitates excessively, and unrecrystallized ferrite having poor ductility remains after annealing, resulting in anisotropy in tensile properties. Increase. Therefore, Ti and V are Ti: 0.100% or less (including 0%) and V: 0.100% or less (including 0%), and Nb, Ti, and V in total are 0.005 to 0.100. %. A preferable total amount is 0.007 to 0.040%.
本発明の高強度鋼板は、以下の成分を任意成分として含有できる。 The high-strength steel sheet of the present invention can contain the following components as optional components.
Cr:0.3%以下
Crは本発明の作用効果を害さない微量元素として含有してもよい。Cr含有量が0.3%を超えると焼入性の向上によりマルテンサイトが過剰に生成して降伏比の低下を招く場合がある。したがって、Crを添加する場合、Cr含有量は0.3%以下とする。Cr: 0.3% or less Cr may be contained as a trace element that does not impair the effects of the present invention. If the Cr content exceeds 0.3%, the martensite may be excessively generated due to the improvement of the hardenability, leading to a decrease in the yield ratio. Therefore, when adding Cr, Cr content shall be 0.3% or less.
Mo:0.3%以下
Moは本発明の作用効果を害さない微量元素として含有してもよい。しかしながら、Mo含有量が0.3%を超えると焼入性の向上によりマルテンサイトが過剰に生成して降伏比の低下を招く場合がある。したがって、Moを添加する場合、Mo含有量は0.3%以下とする。Mo: 0.3% or less Mo may be contained as a trace element that does not impair the effects of the present invention. However, if the Mo content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, which may lead to a decrease in yield ratio. Therefore, when Mo is added, the Mo content is 0.3% or less.
B:0.005%以下
Bは本発明の作用効果を害さない微量元素として含有してもよい。しかしながら、B含有量が0.005%を超えると焼入性の向上によりマルテンサイトが過剰に生成して降伏比の低下を招く場合がある。したがって、Bを添加する場合、B含有量は0.005%以下とする。B: 0.005% or less B may be contained as a trace element that does not impair the effects of the present invention. However, if the B content exceeds 0.005%, martensite may be generated excessively due to the improvement of hardenability, leading to a decrease in yield ratio. Therefore, when adding B, B content shall be 0.005% or less.
Cu:0.3%以下
Cuは本発明の作用効果を害さない微量元素として含有してもよい。しかしながら、Cu含有量が0.3%を超えると焼入性の向上によりマルテンサイトが過剰に生成して降伏比の低下を招く場合がある。したがって、Cuを添加する場合、Cu含有量は0.3%以下とする。Cu: 0.3% or less Cu may be contained as a trace element that does not impair the effects of the present invention. However, if the Cu content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, which may lead to a decrease in yield ratio. Therefore, when adding Cu, Cu content shall be 0.3% or less.
Ni:0.3%以下
Niは本発明の作用効果を害さない微量元素として含有してもよい。しかしながら、Ni含有量が0.3%を超えると焼入性の向上によりマルテンサイトが過剰に生成して降伏比の低下を招く場合がある。したがって、Niを添加する場合、Ni含有量は0.3%以下とする。Ni: 0.3% or less Ni may be contained as a trace element that does not impair the effects of the present invention. However, if the Ni content exceeds 0.3%, martensite may be generated excessively due to the improvement of hardenability, leading to a decrease in yield ratio. Therefore, when Ni is added, the Ni content is 0.3% or less.
Sb:0.3%以下
Sbは本発明の作用効果を害さない微量元素として含有してもよい。しかしながら、Sb含有量が0.3%を超えると高強度鋼板の脆化を招く。したがってSbを添加する場合、Sb含有量は0.3%以下とする。Sb: 0.3% or less Sb may be contained as a trace element that does not impair the effects of the present invention. However, if the Sb content exceeds 0.3%, the high-strength steel sheet is brittle. Therefore, when adding Sb, Sb content shall be 0.3% or less.
上記以外の残部はFeおよび不可避的不純物である。また、本発明では、上記のほかにSn、Co、W、Ca、Na、Mgなどの元素も、本発明の作用効果を害さない微量な範囲で、不可避的不純物として含有してもよい。「微量な範囲」とは、これらの元素を合計で0.01%以下を意味する。 The balance other than the above is Fe and inevitable impurities. In the present invention, in addition to the above, elements such as Sn, Co, W, Ca, Na, and Mg may be contained as inevitable impurities within a minute range that does not impair the effects of the present invention. The “trace range” means 0.01% or less of these elements in total.
(2)鋼組織
本発明の高強度鋼板の鋼組織は、面積率でフェライト:90%以上、パーライトとセメンタイトの合計:0〜10%、マルテンサイトと残留オーステナイトの合計:0〜3%からなる。また、この鋼組織において、上記フェライトの平均結晶粒径は15.0μm以下であり、Nb炭化物とTi炭化物および/またはV炭化物の平均粒子径は5〜50nmであり、Nb炭化物とTi炭化物および/またはV炭化物の析出量の合計は体積率で0.005〜0.050%である。(2) Steel structure The steel structure of the high-strength steel sheet of the present invention is composed of ferrite by area ratio of 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and residual austenite: 0 to 3%. . Further, in this steel structure, the average crystal grain size of the ferrite is 15.0 μm or less, the average particle size of Nb carbide and Ti carbide and / or V carbide is 5 to 50 nm, Nb carbide and Ti carbide and / or Or the total amount of precipitation of V carbide is 0.005 to 0.050% by volume.
フェライト:90%以上
フェライトは良好な延性を有し、鋼組織に主相として含まれ、その含有量は面積率で90%以上である。フェライトの含有量が面積率で90%未満では本発明が目的とする高降伏比が得られず、また、引張特性の異方性も大きくなる。よって、フェライトの含有量は面積率で90%以上とする。好ましくは95%以上とする。なお、本発明の高強度鋼板の鋼組織はフェライト単相(フェライトの含有量が面積率で100%)でもよい。Ferrite: 90% or more Ferrite has good ductility and is contained in the steel structure as a main phase, and its content is 90% or more in terms of area ratio. If the ferrite content is less than 90% by area ratio, the high yield ratio intended by the present invention cannot be obtained, and the anisotropy of tensile properties also increases. Therefore, the ferrite content is 90% or more in terms of area ratio. Preferably, it is 95% or more. The steel structure of the high-strength steel sheet of the present invention may be a ferrite single phase (the ferrite content is 100% in area ratio).
パーライトとセメンタイトの合計:0〜10%
パーライトとセメンタイトは所望の降伏強さと引張強さを得るために有効である。しかし、パーライトとセメンタイトの合計が面積率で10%を超えると本発明が目的とする高降伏比が得られず、引張特性の異方性も大きくなる。このためパーライトとセメンタイトの合計は面積率で0〜10%とする。好ましくは0〜5%とする。Total of pearlite and cementite: 0-10%
Pearlite and cementite are effective in obtaining desired yield strength and tensile strength. However, if the total of pearlite and cementite exceeds 10% in terms of area ratio, the high yield ratio intended by the present invention cannot be obtained, and the anisotropy of tensile properties increases. Therefore, the total of pearlite and cementite is 0 to 10% in terms of area ratio. Preferably it is 0 to 5%.
マルテンサイトと残留オーステナイトの合計:0〜3%
鋼組織は、面積率で、マルテンサイトと残留オーステナイトを合計で0〜3%含有してもよい。マルテンサイトと残留オーステナイトの合計が3%を超えると0.70以上の降伏比が得られなくなる。このためマルテンサイトと残留オーステナイトの合計は0〜3%とする。Total of martensite and retained austenite: 0-3%
The steel structure may contain 0 to 3% of martensite and retained austenite in terms of area ratio. When the sum of martensite and retained austenite exceeds 3%, a yield ratio of 0.70 or more cannot be obtained. Therefore, the total of martensite and retained austenite is 0 to 3%.
フェライトの平均結晶粒径が15.0μm以下
フェライトの平均結晶粒径を所望の範囲に調整することは、本発明が目的とする0.70以上の高降伏比を得るために重要である。フェライトの平均結晶粒径が15.0μmを超えると、0.70以上の降伏比が得られない。したがって、フェライトの平均結晶粒径は15.0μm以下とする。好ましくは10.0μm以下とする。なお、フェライト平均結晶粒径の下限は特に限定されないが、1.0μm未満では引張強さや降伏強さが過度に増加し、曲げ性や伸びの劣化を招く場合があるのでフェライト平均粒径は1.0μm以上であることが好ましい。The average crystal grain size of ferrite is 15.0 μm or less It is important to adjust the average crystal grain size of ferrite to a desired range in order to obtain a high yield ratio of 0.70 or more, which is an object of the present invention. When the average crystal grain size of ferrite exceeds 15.0 μm, a yield ratio of 0.70 or more cannot be obtained. Therefore, the average crystal grain size of ferrite is 15.0 μm or less. Preferably it is 10.0 μm or less. The lower limit of the ferrite average crystal grain size is not particularly limited, but if it is less than 1.0 μm, the tensile strength and yield strength increase excessively, which may lead to deterioration of bendability and elongation. It is preferable that it is 0.0 μm or more.
フェライトの平均アスペクト比が5.0以下
フェライトの平均アスペクト比が5.0を超えると、未再結晶フェライトが増加し、引張特性の異方性が増大する。本発明では、フェライトの平均アスペクト比は4.5以下が好ましく、4.2以下がより好ましい。When the average aspect ratio of ferrite is 5.0 or less, when the average aspect ratio of ferrite exceeds 5.0, non-recrystallized ferrite increases and the anisotropy of tensile properties increases. In the present invention, the average aspect ratio of ferrite is preferably 4.5 or less, and more preferably 4.2 or less.
Nb炭化物とTi炭化物および/またはV炭化物の平均粒子径が5〜50nm
Nb炭化物やTi炭化物やV炭化物は主にフェライト粒内に析出し、その平均粒子径は本発明が目的とする高降伏比と引張特性の等方性を両立するのに重要である。上記粒子径が5nm未満では降伏強さと引張強さが過度に増加するばかりか、引張特性の異方性も増大する。上記粒子径が50nmを超えると降伏強さの増加が不十分となり、本発明が目的とする高降伏比が得られない。よってNb炭化物とTi炭化物および/またはV炭化物の粒子径は5〜50nmとする。下限について好ましい平均粒子径は10nm以上である。上限について好ましい平均粒子径は40nm以下とする。なお、本発明では、Nb炭化物、Ti炭化物、V炭化物を区別せずに平均粒子径を測定する。Average particle size of Nb carbide and Ti carbide and / or V carbide is 5 to 50 nm
Nb carbide, Ti carbide, and V carbide precipitate mainly in the ferrite grains, and the average particle diameter is important for achieving both the high yield ratio and the isotropy of tensile properties aimed by the present invention. If the particle diameter is less than 5 nm, not only the yield strength and tensile strength increase excessively, but also the anisotropy of tensile properties increases. If the particle diameter exceeds 50 nm, the increase in yield strength is insufficient, and the high yield ratio intended by the present invention cannot be obtained. Therefore, the particle diameter of Nb carbide, Ti carbide and / or V carbide is 5 to 50 nm. A preferable average particle size for the lower limit is 10 nm or more. A preferable average particle size for the upper limit is 40 nm or less. In the present invention, the average particle diameter is measured without distinguishing Nb carbide, Ti carbide, and V carbide.
Nb炭化物とTi炭化物およびV炭化物の析出量の合計が体積率で0.005〜0.050%
Nb炭化物とTi炭化物およびV炭化物の析出量を所望の範囲に調整することは、本発明が目的とする高降伏比と引張特性の等方性を両立するのに重要である。Nb炭化物とTi炭化物およびV炭化物の析出量の合計が体積率で0.005%未満だと降伏強さの増加が不十分となり、本発明が目的とする高降伏比が得られない。Nb炭化物とTi炭化物およびV炭化物の析出量の合計が体積率で0.050%を超えるとフェライトの再結晶が顕著に抑制されて降伏強さと引張強さが過度に増加し、さらに、引張特性の異方性が増大する。よってNb炭化物とTi炭化物およびV炭化物の析出量の合計は体積率で0.005〜0.050%とする。下限について好ましい合計体積率は0.010%以上である。上限について好ましい合計体積率は0.040%以下とする。なお、Ti炭化物を含まない場合はTi炭化物を0と考え、V炭化物を含まない場合はV炭化物を0と考える。The total precipitation amount of Nb carbide, Ti carbide and V carbide is 0.005 to 0.050% by volume.
It is important to adjust the precipitation amount of Nb carbide, Ti carbide, and V carbide to a desired range in order to achieve both the high yield ratio and the isotropic property of the tensile properties aimed by the present invention. If the total precipitation amount of Nb carbide, Ti carbide, and V carbide is less than 0.005% by volume, the increase in yield strength is insufficient, and the high yield ratio intended by the present invention cannot be obtained. If the total precipitation amount of Nb carbide, Ti carbide, and V carbide exceeds 0.050% by volume, recrystallization of ferrite is remarkably suppressed, yield strength and tensile strength increase excessively, and tensile properties Increases the anisotropy. Therefore, the total precipitation amount of Nb carbide, Ti carbide, and V carbide is set to 0.005 to 0.050% by volume. A preferable total volume ratio for the lower limit is 0.010% or more. A preferable total volume ratio for the upper limit is 0.040% or less. When Ti carbide is not included, Ti carbide is considered as 0, and when V carbide is not included, V carbide is considered as 0.
なお、各組織の面積率は圧延幅方向に垂直な断面の鋼板表面側から板厚方向に1/4位置を中心とする板厚1/8〜3/8の範囲をSEMで観察し、ASTM E 562−05に記載のポイントカウント法により求める。フェライトの平均結晶粒径は、上記板厚1/4位置を中心とする板厚1/8〜3/8の範囲をSEMで観察し、観察面積と結晶粒数から円相当径を算出することで求める。フェライトの平均アスペクト比は、圧延幅方向に垂直な断面において、鋼板表面から板厚方向に1/4位置を光学顕微鏡で観察し、JIS G 0551の表1に記載の結晶粒当たりの平均線分長を求める方法で圧延方向および板厚方向の平均結晶粒長を算出し、(圧延方向の平均結晶粒長)/(板厚方向の平均結晶粒長)により求めることができる。Nb炭化物やTi炭化物やV炭化物の粒子径は高強度鋼板から薄膜サンプルを作製し、TEM観察像から円相当径を算出(観察面積と粒子数から算出)することにより求める。Nb炭化物とTi炭化物とV炭化物の合計体積率は抽出残渣法により求める。 In addition, the area ratio of each structure was observed by SEM in the range of the plate thickness 1/8 to 3/8 centered on the 1/4 position in the plate thickness direction from the steel plate surface side of the cross section perpendicular to the rolling width direction. It is determined by the point counting method described in E 562-05. For the average crystal grain size of ferrite, the range of the plate thickness 1/8 to 3/8 centered on the above-mentioned plate thickness 1/4 position is observed with SEM, and the equivalent circle diameter is calculated from the observation area and the number of crystal grains. Ask for. The average aspect ratio of the ferrite was measured by observing a quarter position from the steel sheet surface to the plate thickness direction with an optical microscope in a cross section perpendicular to the rolling width direction. The average crystal grain length in the rolling direction and the plate thickness direction is calculated by a method for determining the length, and can be determined by (average crystal grain length in the rolling direction) / (average crystal grain length in the plate thickness direction). The particle diameters of Nb carbide, Ti carbide, and V carbide are obtained by preparing a thin film sample from a high-strength steel plate and calculating the equivalent circle diameter from a TEM observation image (calculating from the observation area and the number of particles). The total volume ratio of Nb carbide, Ti carbide and V carbide is determined by the extraction residue method.
(3)製造条件
本発明の高強度鋼板は上記成分組成を有する鋼を溶製し、鋳造によりスラブ(鋼片)を製造後、熱間圧延、冷間圧延後、連続焼鈍炉で焼鈍を行うことにより製造される。熱間圧延後に酸洗してもよい。以下、熱間圧延工程、冷間圧延工程、焼鈍工程を有する本発明の製造方法について説明する。なお、以下の説明において温度は表面温度を意味する。(3) Manufacturing conditions The high-strength steel sheet of the present invention melts steel having the above component composition, manufactures a slab (steel piece) by casting, performs hot rolling and cold rolling, and then anneals in a continuous annealing furnace. It is manufactured by. Pickling may be performed after hot rolling. Hereinafter, the manufacturing method of this invention which has a hot rolling process, a cold rolling process, and an annealing process is demonstrated. In the following description, temperature means surface temperature.
鋳造方法は特に限定されるものではなく、顕著な成分組成の偏析や組織の不均一が発生しなければ、造塊法、連続鋳造法のいずれで鋳造しても構わない。 The casting method is not particularly limited, and casting may be performed by either the ingot casting method or the continuous casting method as long as the segregation of the significant component composition and the unevenness of the structure do not occur.
熱間圧延は、高温の鋳造スラブをそのまま圧延してもよいし、室温まで冷却されたスラブを再加熱してから圧延してもよい。またスラブの時点で割れなどの表面欠陥がある場合はグラインダーなどによってスラブ手入れを施すことができる。スラブを再加熱する場合は、Nb炭化物とTi炭化物および/またはV炭化物を溶解させるため1100℃以上に加熱することが好ましい。 In hot rolling, a high-temperature cast slab may be rolled as it is, or the slab cooled to room temperature may be reheated and then rolled. If there is a surface defect such as a crack at the time of the slab, the slab can be treated with a grinder. When the slab is reheated, it is preferably heated to 1100 ° C. or higher in order to dissolve Nb carbide, Ti carbide and / or V carbide.
熱間圧延工程とは、鋼を熱間圧延し、該熱間圧延後、仕上圧延温度〜650℃の温度域の滞留時間を10秒以下の条件で鋼板を冷却し、500〜700℃で巻取る工程である。 In the hot rolling process, steel is hot rolled, and after the hot rolling, the steel sheet is cooled under a condition that the residence time in the temperature range of finish rolling temperature to 650 ° C. is 10 seconds or less, and wound at 500 to 700 ° C. It is a process to take.
熱間圧延では、スラブに粗圧延、仕上圧延を施す。その後、熱間圧延後の鋼板を巻取り熱延コイルとする。熱間圧延における粗圧延条件および仕上圧延条件は特に限定されるものではなく常法にしたがって決定すればよい。仕上圧延温度がAr3点未満になると、熱延鋼板の鋼組織中に圧延方向に伸長した粗大なフェライトが生成し、焼鈍後に延性の低下を招く場合がある。このため、仕上圧延温度はAr3点以上とすることが好ましい。なお、Ar3点は変態点測定装置(例えばフォーマスター試験機)を用いてオーステナイト単相温度域から1℃/sで連続冷却したときにフェライト変態が開始する温度を測定することで求めることができる。 In hot rolling, the slab is subjected to rough rolling and finish rolling. Then, the hot-rolled steel sheet is taken up as a hot rolled coil. The rough rolling conditions and finish rolling conditions in the hot rolling are not particularly limited, and may be determined according to a conventional method. When the finish rolling temperature is less than the Ar3 point, coarse ferrite stretched in the rolling direction is generated in the steel structure of the hot-rolled steel sheet, and ductility may be lowered after annealing. For this reason, it is preferable that finishing rolling temperature shall be Ar3 point or more. In addition, Ar3 point can be calculated | required by measuring the temperature which a ferrite transformation starts when it cools continuously at 1 degree-C / s from an austenite single phase temperature range using a transformation point measuring apparatus (for example, for master test machine). .
仕上圧延温度〜650℃の温度域の滞留時間:10秒以下
熱間圧延後の冷却において、仕上圧延温度〜650℃の温度域の滞留時間を適正に制御することで、フェライトの平均結晶粒径の粗大化を抑制することができる。このため、上記冷却条件は、本発明において重要である。仕上圧延後の冷却において仕上圧延温度〜650℃の温度域の滞留時間が10秒を超えると、熱間圧延の巻取後に粗大なNb炭化物やTi炭化物やV炭化物が過度に析出するため、焼鈍時にフェライト粒が粗大になりやすくなりフェライトの平均結晶粒径が15.0μmを超えるため降伏比が低下する。そこで、上記冷却における、仕上圧延温度〜650℃の温度域の滞留時間は10秒以下とする。なお、上記滞留時間の下限は特に限定されないが、焼鈍時に均一にNb炭化物やTi炭化物やV炭化物を析出させてフェライト結晶粒径を均一にする観点から1秒以上滞留することが好ましい。また、上記滞留時間が制御される温度域の下限は、Nb炭化物等の平均粒子径が本発明範囲外となったり、Nb炭化物等の析出量の合計が本発明範囲外となるのを抑えるという理由で650℃とする。Retention time in the temperature range from the finish rolling temperature to 650 ° C .: 10 seconds or less In cooling after hot rolling, the average grain size of ferrite is controlled by appropriately controlling the residence time in the temperature range from the finish rolling temperature to 650 ° C. Can be suppressed. For this reason, the said cooling conditions are important in this invention. If the residence time in the temperature range of the finish rolling temperature to 650 ° C. exceeds 10 seconds in the cooling after finish rolling, coarse Nb carbide, Ti carbide and V carbide are excessively precipitated after hot rolling, so annealing is performed. At times, the ferrite grains tend to be coarse and the average crystal grain size of ferrite exceeds 15.0 μm, so the yield ratio decreases. Therefore, the residence time in the temperature range of the finish rolling temperature to 650 ° C. in the cooling is 10 seconds or less. The lower limit of the residence time is not particularly limited, but it is preferably retained for 1 second or more from the viewpoint of uniformly depositing Nb carbide, Ti carbide, and V carbide during annealing to make the ferrite crystal grain size uniform. Further, the lower limit of the temperature range in which the residence time is controlled is to suppress that the average particle diameter of Nb carbide and the like is outside the range of the present invention, and the total amount of precipitation of Nb carbide and the like is outside the range of the present invention. It is set to 650 ° C. for the reason.
巻取温度:500〜700℃
巻取温度は、Nb炭化物やTi炭化物やV炭化物の析出量およびこれらの平均粒子径の調整により、焼鈍後のフェライト平均結晶粒径を15.0μm以下に制御するために重要である。鋼板の幅方向中央において、巻取温度が500℃未満では巻取後の冷却中に上記炭化物が十分析出せず、焼鈍の加熱および均熱時に粗大な炭化物が析出し、フェライト粒径が粗大化するため、高降伏比が得られず、さらに引張強さも小さくなる。巻取温度が700℃を超えると巻取後の冷却中に粗大なNb炭化物やTi炭化物やV炭化物が析出し、焼鈍時にフェライト粒径が粗大化するため、高降伏比が得られず、さらに引張強さも小さくなる。したがって巻取温度は500〜700℃とする。下限について好ましい巻取温度は550℃以上である。上限について好ましい巻取温度は650℃以下である。Winding temperature: 500-700 ° C
The coiling temperature is important for controlling the ferrite average crystal grain size after annealing to 15.0 μm or less by adjusting the precipitation amount of Nb carbide, Ti carbide, V carbide and the average particle diameter thereof. When the coiling temperature is less than 500 ° C at the center in the width direction of the steel sheet, the carbides do not sufficiently precipitate during cooling after winding, coarse carbides precipitate during annealing and soaking, and the ferrite grain size increases. Therefore, a high yield ratio cannot be obtained, and the tensile strength is also reduced. When the coiling temperature exceeds 700 ° C., coarse Nb carbide, Ti carbide and V carbide precipitate during cooling after winding, and the ferrite grain size becomes coarse during annealing, so a high yield ratio cannot be obtained. The tensile strength is also reduced. Therefore, the coiling temperature is 500 to 700 ° C. A preferable winding temperature for the lower limit is 550 ° C. or higher. A preferable coiling temperature for the upper limit is 650 ° C. or lower.
冷間圧延工程とは、上記熱間圧延工程で得られた熱延鋼板を冷間圧延する工程である。冷間圧延の圧延率は75%以下とする。好ましくは30〜75%である。圧延率が75%を超えると炭化物の平均粒子径が粗大になり本願が目的とする高YRが得られないため75%以下が必要である。圧延率が30%以上であれば焼鈍時にフェライトを完全に再結晶させて等方的な引張特性が得られるため好ましい。 The cold rolling step is a step of cold rolling the hot rolled steel sheet obtained in the hot rolling step. The rolling rate of cold rolling is 75% or less. Preferably it is 30 to 75%. If the rolling rate exceeds 75%, the average particle diameter of the carbide becomes coarse, and the high YR intended by the present application cannot be obtained, so 75% or less is necessary. A rolling rate of 30% or more is preferable because isotropic tensile properties can be obtained by completely recrystallizing ferrite during annealing.
焼鈍は、連続焼鈍炉を用いて、均熱温度まで昇温後、冷却する工程からなる。本発明における焼鈍工程とは、冷間圧延工程で得られる冷延鋼板を、連続焼鈍炉で、昇温時における650〜750℃の温度域で滞留時間:60秒以下で滞留し、該滞留後に均熱温度:760〜880℃、均熱時間;120秒以下の条件で均熱し、該均熱後400〜500℃の温度域の滞留時間が100秒以下の条件で冷却する工程である。 Annealing consists of a process of raising the temperature to a soaking temperature and then cooling it using a continuous annealing furnace. The annealing process in the present invention means that the cold-rolled steel sheet obtained in the cold rolling process is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at the time of temperature rise in a residence time of 60 seconds or less, and after the residence Soaking temperature: 760 to 880 ° C., soaking time; soaking under conditions of 120 seconds or less, and after the soaking, cooling is performed under conditions of a residence time in a temperature range of 400 to 500 ° C. of 100 seconds or less.
昇温時に650〜750℃の温度域での滞留時間:60秒以下
昇温時の650〜750℃における滞留時間は焼鈍後のフェライトの平均アスペクト比を5.0以下に制御するために重要な製造条件である。昇温時の650〜750℃における滞留時間が60秒を超えると、フェライトが圧延方向に粒成長しやすくなるためフェライトの平均アスペクト比が5.0を超える。よって昇温時の650〜750℃における滞留時間は60秒以下とする。好ましくは昇温時の650〜750℃における滞留時間は50秒以下とする。なお、滞留時間の下限は特に限定されないが、滞留時間が短すぎるとフェライトの再結晶が十分に進行しないため、滞留時間は5秒以上が好ましい。Residence time in the temperature range of 650 to 750 ° C. at the time of temperature rise: 60 seconds or less The residence time at 650 to 750 ° C. at the time of temperature rise is important for controlling the average aspect ratio of the ferrite after annealing to 5.0 or less. Manufacturing conditions. When the residence time at 650 to 750 ° C. at the time of temperature rise exceeds 60 seconds, the ferrite easily grows in the rolling direction, so the average aspect ratio of ferrite exceeds 5.0. Therefore, the residence time at 650 to 750 ° C. during the temperature rise is set to 60 seconds or less. Preferably, the residence time at 650 to 750 ° C. when the temperature is raised is 50 seconds or less. The lower limit of the residence time is not particularly limited, but if the residence time is too short, the recrystallization of ferrite does not proceed sufficiently, so the residence time is preferably 5 seconds or more.
均熱温度:760〜880℃、均熱時間:120秒以下
均熱温度および均熱時間はフェライト平均結晶粒径を制御する上で重要な条件である。均熱温度が760℃未満ではフェライトの再結晶が不十分となり引張特性の異方性が増大する。均熱温度が880℃を超えるとフェライト平均結晶粒径が粗大化して本発明が目的とする降伏比が得られず、引張強さも小さくなる。このため均熱温度は760〜880℃とする。また均熱時間が120秒を超えると、フェライト平均結晶粒径が粗大化するため本発明が目的とする引張強さと高降伏比が得られない。このため均熱時間は120秒以下とする。好ましくは60秒以下とする。なお、均熱時間の下限は特に限定されないが、引張特性の異方性低減の観点からフェライトを完全に再結晶させることが好ましいため均熱時間は30秒以上が好ましい。Soaking temperature: 760 to 880 ° C., soaking time: 120 seconds or less Soaking temperature and soaking time are important conditions for controlling the ferrite average crystal grain size. If the soaking temperature is less than 760 ° C., the recrystallization of ferrite becomes insufficient and the anisotropy of tensile properties increases. When the soaking temperature exceeds 880 ° C., the ferrite average crystal grain size becomes coarse, the desired yield ratio of the present invention cannot be obtained, and the tensile strength is also reduced. Therefore, the soaking temperature is set to 760 to 880 ° C. On the other hand, if the soaking time exceeds 120 seconds, the average grain size of ferrite becomes coarse, so that the intended tensile strength and high yield ratio of the present invention cannot be obtained. For this reason, the soaking time is 120 seconds or less. Preferably it is 60 seconds or less. The lower limit of the soaking time is not particularly limited, but from the viewpoint of reducing the anisotropy of the tensile properties, it is preferable to completely recrystallize the ferrite, so the soaking time is preferably 30 seconds or more.
昇温および均熱時の加熱方式は特に限定されるものではなく、ラジアントチューブ方式や直火加熱方式などで行うことができる。 The heating method at the time of temperature increase and soaking is not particularly limited, and can be performed by a radiant tube method, a direct fire heating method, or the like.
均熱後の冷却における冷却条件は、400〜500℃の温度域の滞留時間が100秒以下である。滞留時間が100秒以下であることは炭化物の平均粒子径を50nm以下にするために必要である。なお滞留時間の下限は特に限定されないが、極端に短くするとフェライト中の固溶Cが増加して耐時効特性が劣化したり、冷却設備への過度な投資が必要になるため5秒以上が好ましい。より好ましくは10秒以上である。ここで、「400〜500℃の温度域の滞留時間」とは、冷却中の鋼板が400〜500℃の温度になっている時間の合計を意味し、冷却停止温度が400℃以上であれば、冷却停止温度から500℃になっている時間の合計を意味する。また、この温度域での滞留は過時効処理に相当する。なお、その他の冷却条件は特に限定されないが、冷却停止温度が400〜500℃、平均冷却速度30℃/s以下の条件が挙げられる。 The cooling condition in the cooling after soaking is that the residence time in the temperature range of 400 to 500 ° C. is 100 seconds or less. It is necessary for the residence time to be 100 seconds or less to make the average particle size of the carbides 50 nm or less. The lower limit of the residence time is not particularly limited. However, if it is extremely short, the solid solution C in the ferrite is increased and the aging resistance is deteriorated or excessive investment in the cooling equipment is required. . More preferably, it is 10 seconds or more. Here, “the residence time in the temperature range of 400 to 500 ° C.” means the total time during which the steel plate being cooled is at a temperature of 400 to 500 ° C. If the cooling stop temperature is 400 ° C. or higher. Means the total time from the cooling stop temperature to 500 ° C. Moreover, the residence in this temperature range is equivalent to an overaging treatment. In addition, although other cooling conditions are not specifically limited, Cooling stop temperature is 400-500 degreeC, and the conditions whose average cooling rate is 30 degrees C / s or less are mentioned.
上記のようにして得られた高強度鋼板の表面にめっきを施すことができる。めっきは亜鉛めっきが好ましく、本発明の高強度鋼板に亜鉛めっきを施すことで、高強度鋼板上に亜鉛めっき層が形成される。亜鉛めっき(電気亜鉛めっき、溶融亜鉛めっき等)の中でも、溶融亜鉛めっき浴に浸漬する溶融亜鉛めっきが好適である。 Plating can be applied to the surface of the high-strength steel plate obtained as described above. The plating is preferably galvanization, and the galvanization layer is formed on the high-strength steel plate by applying galvanization to the high-strength steel plate of the present invention. Among galvanizing (electrogalvanizing, hot dip galvanizing, etc.), hot dip galvanizing immersed in a hot dip galvanizing bath is suitable.
高強度鋼板に溶融亜鉛めっきを施すことで形成される溶融亜鉛めっき層に対して合金化処理を施すことにより、合金化溶融亜鉛めっき層が形成される。合金化処理を施す場合、保持温度が450℃未満では十分に合金化が進まずめっき密着性や耐食性が劣化する場合がある。また、保持温度が560℃を超えると合金化が過度に進行してプレス時にパウダリングなどの問題が発生する場合がある。このため保持温度は450〜560℃とするのが好ましい。また、保持時間が5秒未満では十分に合金化が進まずめっき密着性や耐食性が劣化する場合があるため、保持時間は5秒以上とすることが好ましい。 An alloyed hot-dip galvanized layer is formed by subjecting a hot-dip galvanized layer formed by hot-dip galvanizing to a high-strength steel plate. When the alloying treatment is performed, if the holding temperature is less than 450 ° C., alloying does not proceed sufficiently, and the plating adhesion and corrosion resistance may deteriorate. Further, when the holding temperature exceeds 560 ° C., alloying proceeds excessively, and problems such as powdering may occur during pressing. Therefore, the holding temperature is preferably 450 to 560 ° C. Further, if the holding time is less than 5 seconds, alloying does not proceed sufficiently and the plating adhesion and corrosion resistance may deteriorate, so the holding time is preferably 5 seconds or more.
その後、必要に応じて伸長率0.1〜5.0%の調質圧延を施してもよい。 Thereafter, temper rolling with an elongation of 0.1 to 5.0% may be performed as necessary.
以上により、本発明の目的とする高強度鋼板が得られる。本発明の高強度鋼板に対して、化成処理、有機系皮膜処理等の表面処理、塗装を施しても本発明の目的とする特性を損なうことはない。 As described above, the high-strength steel plate targeted by the present invention is obtained. Even if the high-strength steel sheet of the present invention is subjected to a surface treatment such as a chemical conversion treatment or an organic coating treatment or a coating, the target properties of the present invention are not impaired.
以下、本発明を実施例により詳細に説明する。 Hereinafter, the present invention will be described in detail with reference to examples.
表1に示す成分組成を有する鋼A〜Mの鋼スラブを1250℃で1時間均熱後、仕上板厚3.2mm、Ar3点以上である仕上圧延温度900℃の条件で圧延後、表2に示す条件で冷却し、表2に示す巻取温度で巻取った。製造した熱延鋼板を酸洗後、仕上板厚1.4mmの冷間圧延を施して冷延鋼板とし、表2に示す条件の焼鈍を施してNo.1〜31の高強度鋼板を製造した。なお、焼鈍における冷却の冷却条件は、冷却停止温度が480℃、平均冷却速度20℃/s以下、400〜500℃の温度域(500℃〜冷却停止温度の温度域)での滞留時間30秒とした。焼鈍は、めっきを施さない場合はCALを用いて行った。また、めっきを施す場合はCGLを用いて、溶融亜鉛めっきまたは合金化溶融亜鉛めっきを施した。めっき層を合金化溶融亜鉛めっき層とする場合は、510℃で10秒保持する合金化処理を施した。 After steel slabs of steels A to M having the composition shown in Table 1 were soaked at 1250 ° C. for 1 hour, after rolling at a finish sheet thickness of 3.2 mm and a finish rolling temperature of 900 ° C. that is Ar 3 or higher, Table 2 The sample was cooled under the conditions shown in FIG. The picked hot-rolled steel sheet was pickled and then cold-rolled to a finished sheet thickness of 1.4 mm to obtain a cold-rolled steel sheet, and annealed under the conditions shown in Table 2 to obtain No. 1 to 31 high strength steel sheets were produced. In addition, the cooling conditions for cooling in annealing are as follows: the cooling stop temperature is 480 ° C., the average cooling rate is 20 ° C./s or less, and the residence time in the temperature range of 400 to 500 ° C. (500 ° C. to the temperature range of cooling stop temperature) is 30 seconds. It was. Annealing was performed using CAL when plating was not performed. Moreover, when performing plating, CGL was used, and hot dip galvanization or galvannealing was performed. When the plated layer was an alloyed hot-dip galvanized layer, an alloying treatment was performed by holding at 510 ° C. for 10 seconds.
得られた高強度鋼板に対し、鋼組織観察と引張試験を行った。 Steel structure observation and a tensile test were performed on the obtained high-strength steel sheet.
鋼組織の面積率は、各組織の面積率は圧延幅方向に垂直な断面の鋼板表面側から板厚方向に1/4位置を中心とする板厚1/8〜3/8の範囲をSEMで観察し、ASTM E 562−05に記載のポイントカウント法により求めた。フェライトの平均結晶粒径は、上記板厚1/4位置を中心とする板厚1/8〜3/8の範囲をSEMで観察し、観察面積と結晶粒数から円相当径を算出することで求めた。フェライトの平均アスペクト比は、圧延幅方向に垂直な断面において、鋼板表面から板厚方向に1/4位置を光学顕微鏡で観察し、JIS G 0551の表1に記載の結晶粒当たりの平均線分長を求める方法で圧延方向および板厚方向の平均結晶粒長を算出し、(圧延方向の平均結晶粒長)/(板厚方向の平均結晶粒長)により求めた。炭化物(Nb炭化物、Ti炭化物、V炭化物)の平均粒子径はTEM観察を行い、画像処理により円相当径を求めた。Nb炭化物とTi炭化物とV炭化物の合計体積率は抽出残渣法により求めた。観察はいずれも各10視野で行い、その平均を算出した。なお、結果は表2(表2−1と表2−2を合わせて表2とする。)に示し、表2のαがフェライト、Pがパーライト、Mがマルテンサイト、θがセメンタイトを意味し、α粒径がフェライト平均結晶粒径を意味し、M(C、N)粒子径が炭化物の平均粒子径、M(C、N)体積率がNb炭化物とTi炭化物とV炭化物の析出量の合計を意味する。また、上記M(C、N)におけるMは、Nb、Ti又はVを意味する。 As for the area ratio of each steel structure, the area ratio of each structure is an SEM in the range of 1/8 to 3/8 of the sheet thickness centered at 1/4 position in the sheet thickness direction from the steel sheet surface side of the cross section perpendicular to the rolling width direction. And determined by the point counting method described in ASTM E 562-05. For the average crystal grain size of ferrite, the range of the plate thickness 1/8 to 3/8 centered on the above-mentioned plate thickness 1/4 position is observed with SEM, and the equivalent circle diameter is calculated from the observation area and the number of crystal grains. I asked for it. The average aspect ratio of the ferrite was measured by observing a quarter position from the steel sheet surface in the sheet thickness direction with an optical microscope in a cross section perpendicular to the rolling width direction, and the average line segment per crystal grain described in Table 1 of JIS G 0551 The average crystal grain length in the rolling direction and the plate thickness direction was calculated by a method for determining the length, and the average crystal grain length in the rolling direction / (average crystal grain length in the plate thickness direction) was calculated. The average particle diameter of the carbides (Nb carbide, Ti carbide, V carbide) was observed by TEM, and the equivalent circle diameter was determined by image processing. The total volume fraction of Nb carbide, Ti carbide and V carbide was determined by the extraction residue method. All observations were made with 10 fields of view, and the average was calculated. The results are shown in Table 2 (Table 2-1 and Table 2-2 are combined into Table 2). In Table 2, α represents ferrite, P represents pearlite, M represents martensite, and θ represents cementite. , Α grain size means ferrite average crystal grain size, M (C, N) particle diameter is the average particle diameter of carbide, M (C, N) volume ratio is the precipitation amount of Nb carbide, Ti carbide and V carbide Mean total. M in the above M (C, N) means Nb, Ti or V.
引張強さ(TS)および降伏比(YR)は、引張方向が圧延方向と垂直になるよう採取したJIS5号引張試験片を用いて、JIS Z 2241に準拠した引張試験により求めた。引張特性の異方性は、引張方向が圧延方向と平行な向きで行った引張試験における降伏強さと、引張方向が圧延方向と垂直な向きで行った引張試験における降伏強さの差で評価し、差が30MPa以下のものを「○」、30MPaを超えるものを「×」とした。なお、330MPa〜500MPa未満の引張強さ、0.70以上の降伏比を良好と評価した。 The tensile strength (TS) and the yield ratio (YR) were determined by a tensile test in accordance with JIS Z 2241 using a JIS No. 5 tensile test specimen collected so that the tensile direction was perpendicular to the rolling direction. The anisotropy of tensile properties is evaluated by the difference between the yield strength in a tensile test conducted with the tensile direction parallel to the rolling direction and the yield strength in the tensile test conducted with the tensile direction perpendicular to the rolling direction. A sample having a difference of 30 MPa or less was designated as “◯”, and a sample having a difference exceeding 30 MPa was designated as “x”. In addition, the tensile strength of 330 MPa-less than 500 MPa and the yield ratio of 0.70 or more were evaluated as favorable.
表2に鋼組織の観察結果、引張試験結果を示す。No.1〜3、6、8、9、12〜16、18、19、22、24、25、28は本発明の要件をすべて満たしているため、本発明が目的とする高降伏比で引張特性の異方性が小さい高強度鋼板が得られている。一方、No.4、5、7、10、11、17、20、21、23、26、27、29、30、31は成分組成もしくは製造条件が本発明の範囲外であり、所望の鋼組織が得られていないため本発明が目的とする高強度鋼板が得られていない。 Table 2 shows the observation results of the steel structure and the tensile test results. No. 1 to 3, 6, 8, 9, 12 to 16, 18, 19, 22, 24, 25, and 28 all satisfy the requirements of the present invention, so that the tensile properties at the high yield ratio intended by the present invention are high. A high-strength steel sheet with small anisotropy is obtained. On the other hand, no. 4, 5, 7, 10, 11, 17, 20, 21, 23, 26, 27, 29, 30, 31 have component compositions or production conditions outside the scope of the present invention, and a desired steel structure is obtained. Therefore, the high strength steel plate intended by the present invention is not obtained.
本発明の高強度鋼板は、自動車内板部品などを中心に、高降伏比と引張特性の等方性が要求される分野に好適である。
The high-strength steel sheet of the present invention is suitable for a field that requires high yield ratio and isotropy of tensile properties, mainly for automobile inner plate parts.
Claims (11)
鋼組織は、面積率でフェライト:90%以上、パーライトとセメンタイトの合計:0〜10%、マルテンサイトと残留オーステナイトの合計:0〜3%からなり、
前記フェライトの平均結晶粒径が15.0μm以下、平均アスペクト比が5.0以下であり、
Nb炭化物並びにTi炭化物および/またはV炭化物を含み、該Nb炭化物並びにTi炭化物および/またはV炭化物の平均粒子径が5〜50nmであり、
Nb炭化物、Ti炭化物およびV炭化物の析出量の合計が体積率で0.005〜0.050%である高強度鋼板。Component composition is mass%, C: 0.02% to less than 0.10%, Si: less than 0.10%, Mn: less than 1.0%, P: 0.10% or less, S: 0.020 %: Al: 0.01-0.10%, N: 0.010% or less, Nb: 0.005-0.070%, Ti: 0.100% or less (including 0%), V: 0 .100% or less (including 0%) and Nb, Ti and V in total 0.005 to 0.100%, with the balance being Fe and inevitable impurities,
Steel structure consists of ferrite in area ratio: 90% or more, total of pearlite and cementite: 0 to 10%, total of martensite and retained austenite: 0 to 3%,
The average crystal grain size of the ferrite is 15.0 μm or less, the average aspect ratio is 5.0 or less,
Nb carbide and Ti carbide and / or V carbide, the Nb carbide and Ti carbide and / or V carbide has an average particle size of 5 to 50 nm,
A high-strength steel sheet in which the total precipitation amount of Nb carbide, Ti carbide and V carbide is 0.005 to 0.050% by volume.
鋼を熱間圧延し、該熱間圧延後、仕上圧延温度〜650℃の温度域の滞留時間を10秒以下の条件で鋼板を冷却し、500〜700℃で巻取る熱間圧延工程と、
前記熱間圧延工程で得られる熱延鋼板を75%以下の圧延率で冷間圧延する冷間圧延工程と、
前記冷間圧延工程で得られる冷延鋼板を、連続焼鈍炉で、昇温時における650〜750℃の温度域で滞留時間:60秒以下で滞留し、該滞留後に均熱温度:760〜880℃、均熱時間:120秒以下の条件で均熱し、400〜500℃の温度域の滞留時間が100秒以下の条件で冷却する焼鈍工程と、を有する高強度鋼板の製造方法。It is a manufacturing method of the high strength steel plate according to claim 1 or 2,
Hot rolling the steel, and after the hot rolling, a hot rolling step of cooling the steel plate under a condition that the residence time in the temperature range of the finish rolling temperature to 650 ° C. is 10 seconds or less and winding at 500 to 700 ° C .;
A cold rolling step of cold rolling the hot rolled steel sheet obtained in the hot rolling step at a rolling rate of 75% or less;
The cold-rolled steel sheet obtained in the cold rolling step is retained in a continuous annealing furnace in a temperature range of 650 to 750 ° C. at the time of temperature increase in a residence time: 60 seconds or less, and after the residence, a soaking temperature: 760 to 880. C., soaking time: An annealing step of soaking under conditions of 120 seconds or less and cooling under conditions where the residence time in the temperature range of 400 to 500.degree. C. is 100 seconds or less.
The method for producing a high-strength steel sheet according to claim 8, wherein the plating process is an electrogalvanizing process.
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