WO2011090180A1 - High-strength hot-dip galvanized steel sheet with excellent material stability and processability and process for producing same - Google Patents
High-strength hot-dip galvanized steel sheet with excellent material stability and processability and process for producing same Download PDFInfo
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- WO2011090180A1 WO2011090180A1 PCT/JP2011/051151 JP2011051151W WO2011090180A1 WO 2011090180 A1 WO2011090180 A1 WO 2011090180A1 JP 2011051151 W JP2011051151 W JP 2011051151W WO 2011090180 A1 WO2011090180 A1 WO 2011090180A1
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- steel sheet
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- galvanized steel
- dip galvanized
- material stability
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- 229910001335 Galvanized steel Inorganic materials 0.000 title claims abstract description 47
- 239000008397 galvanized steel Substances 0.000 title claims abstract description 47
- 239000000463 material Substances 0.000 title claims abstract description 42
- 238000000034 method Methods 0.000 title abstract description 8
- 229910000831 Steel Inorganic materials 0.000 claims abstract description 73
- 239000010959 steel Substances 0.000 claims abstract description 73
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 54
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 45
- 229910001562 pearlite Inorganic materials 0.000 claims abstract description 29
- 239000000203 mixture Substances 0.000 claims abstract description 23
- 239000012535 impurity Substances 0.000 claims abstract description 6
- 229910052748 manganese Inorganic materials 0.000 claims abstract description 6
- 229910052757 nitrogen Inorganic materials 0.000 claims abstract description 5
- 229910052782 aluminium Inorganic materials 0.000 claims abstract description 4
- 229910052698 phosphorus Inorganic materials 0.000 claims abstract description 3
- 238000011282 treatment Methods 0.000 claims description 16
- 238000005275 alloying Methods 0.000 claims description 15
- 238000005246 galvanizing Methods 0.000 claims description 15
- 238000004519 manufacturing process Methods 0.000 claims description 15
- 238000007747 plating Methods 0.000 claims description 12
- 229910052758 niobium Inorganic materials 0.000 claims description 7
- 229910052720 vanadium Inorganic materials 0.000 claims description 6
- 229910052802 copper Inorganic materials 0.000 claims description 5
- 229910052750 molybdenum Inorganic materials 0.000 claims description 5
- 229910052759 nickel Inorganic materials 0.000 claims description 5
- 229910052718 tin Inorganic materials 0.000 claims description 3
- HCHKCACWOHOZIP-UHFFFAOYSA-N Zinc Chemical compound [Zn] HCHKCACWOHOZIP-UHFFFAOYSA-N 0.000 claims description 2
- 229910052804 chromium Inorganic materials 0.000 claims description 2
- 229910052725 zinc Inorganic materials 0.000 claims description 2
- 239000011701 zinc Substances 0.000 claims description 2
- 229910052799 carbon Inorganic materials 0.000 abstract description 4
- 229910052742 iron Inorganic materials 0.000 abstract description 2
- 229910001566 austenite Inorganic materials 0.000 description 20
- 238000000137 annealing Methods 0.000 description 15
- 238000005728 strengthening Methods 0.000 description 12
- 230000000694 effects Effects 0.000 description 11
- 238000001816 cooling Methods 0.000 description 9
- 238000010438 heat treatment Methods 0.000 description 9
- 230000003647 oxidation Effects 0.000 description 9
- 238000007254 oxidation reaction Methods 0.000 description 9
- 230000000717 retained effect Effects 0.000 description 9
- 230000007423 decrease Effects 0.000 description 7
- 238000010586 diagram Methods 0.000 description 6
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 6
- 239000010410 layer Substances 0.000 description 6
- 229910001035 Soft ferrite Inorganic materials 0.000 description 5
- 230000000052 comparative effect Effects 0.000 description 5
- 239000002131 composite material Substances 0.000 description 5
- 230000006866 deterioration Effects 0.000 description 5
- 238000005098 hot rolling Methods 0.000 description 5
- 238000005121 nitriding Methods 0.000 description 5
- 230000032683 aging Effects 0.000 description 4
- 229910045601 alloy Inorganic materials 0.000 description 4
- 239000000956 alloy Substances 0.000 description 4
- 239000010960 cold rolled steel Substances 0.000 description 4
- 238000005097 cold rolling Methods 0.000 description 4
- 238000005554 pickling Methods 0.000 description 4
- 238000005096 rolling process Methods 0.000 description 4
- 230000009466 transformation Effects 0.000 description 4
- 229910001563 bainite Inorganic materials 0.000 description 3
- 229910001567 cementite Inorganic materials 0.000 description 3
- 239000000446 fuel Substances 0.000 description 3
- 238000007654 immersion Methods 0.000 description 3
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 3
- 150000001247 metal acetylides Chemical class 0.000 description 3
- 239000002244 precipitate Substances 0.000 description 3
- 239000006104 solid solution Substances 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- 239000002344 surface layer Substances 0.000 description 3
- 238000005266 casting Methods 0.000 description 2
- 238000009749 continuous casting Methods 0.000 description 2
- 238000005261 decarburization Methods 0.000 description 2
- 230000003247 decreasing effect Effects 0.000 description 2
- 230000007547 defect Effects 0.000 description 2
- 238000011161 development Methods 0.000 description 2
- 230000018109 developmental process Effects 0.000 description 2
- 238000005265 energy consumption Methods 0.000 description 2
- 238000007710 freezing Methods 0.000 description 2
- 230000008014 freezing Effects 0.000 description 2
- 230000001771 impaired effect Effects 0.000 description 2
- 230000000977 initiatory effect Effects 0.000 description 2
- 239000002184 metal Substances 0.000 description 2
- 229910052751 metal Inorganic materials 0.000 description 2
- 238000005498 polishing Methods 0.000 description 2
- 238000001556 precipitation Methods 0.000 description 2
- 230000001737 promoting effect Effects 0.000 description 2
- 238000005482 strain hardening Methods 0.000 description 2
- 229910052715 tantalum Inorganic materials 0.000 description 2
- 229910018072 Al 2 O 3 Inorganic materials 0.000 description 1
- 241000219307 Atriplex rosea Species 0.000 description 1
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 1
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 1
- 230000002411 adverse Effects 0.000 description 1
- 230000015572 biosynthetic process Effects 0.000 description 1
- 230000001276 controlling effect Effects 0.000 description 1
- 238000012937 correction Methods 0.000 description 1
- 238000005336 cracking Methods 0.000 description 1
- 230000003111 delayed effect Effects 0.000 description 1
- 238000013461 design Methods 0.000 description 1
- 230000005611 electricity Effects 0.000 description 1
- 230000007613 environmental effect Effects 0.000 description 1
- 238000011156 evaluation Methods 0.000 description 1
- 229910052738 indium Inorganic materials 0.000 description 1
- 238000000465 moulding Methods 0.000 description 1
- 238000012545 processing Methods 0.000 description 1
- 238000001953 recrystallisation Methods 0.000 description 1
- 230000001105 regulatory effect Effects 0.000 description 1
- 230000002040 relaxant effect Effects 0.000 description 1
- -1 retained austenite Chemical class 0.000 description 1
- 229920006395 saturated elastomer Polymers 0.000 description 1
- 238000005204 segregation Methods 0.000 description 1
- 238000002791 soaking Methods 0.000 description 1
- 230000006641 stabilisation Effects 0.000 description 1
- 238000011105 stabilization Methods 0.000 description 1
- 230000000087 stabilizing effect Effects 0.000 description 1
- 238000009628 steelmaking Methods 0.000 description 1
- 238000005496 tempering Methods 0.000 description 1
- 238000009864 tensile test Methods 0.000 description 1
- 238000012360 testing method Methods 0.000 description 1
- 150000003568 thioethers Chemical class 0.000 description 1
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Images
Classifications
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
- C21D8/0405—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing of ferrous alloys
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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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
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—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
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
-
- 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
- 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
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
- C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
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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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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/26—After-treatment
- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
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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
- C21D2211/009—Pearlite
Definitions
- the present invention relates to a high-strength hot-dip galvanized steel sheet excellent in material stability and workability suitable as a member used in industrial fields such as automobiles and electricity, and a method for producing the same.
- the shape freezing property is significantly reduced by increasing the strength and thinning of the steel sheet.
- it is widely performed to predict the shape change after mold release in advance during press forming and to design the mold in consideration of the amount of shape change, but when the tensile strength (TS) of the steel sheet changes.
- TS tensile strength
- Patent Document 1 proposes a steel sheet having excellent ductility by specifying chemical components, volume ratios of retained austenite and martensite, and manufacturing methods thereof.
- the steel plate excellent in ductility is proposed by prescribing
- Patent Document 3 proposes a steel sheet having excellent ductility by defining chemical components and defining volume fractions of ferrite, bainitic ferrite and retained austenite.
- Patent Document 4 proposes a method for manufacturing a high-strength cold-rolled steel sheet in which variation in elongation in the sheet width direction is improved.
- Patent Documents 1 to 3 the main purpose is to improve the ductility of the high-strength thin steel sheet, and therefore the hole expandability is not considered.
- Patent Document 4 describes only the variation of the total elongation EL in the plate width direction, and does not consider the variation of the material due to the component composition and manufacturing conditions. Therefore, the development of a high-strength hot-dip galvanized steel sheet having both high ductility and high hole expansibility and excellent material stability becomes an issue.
- the inventors have made extensive studies to obtain a high-strength hot-dip galvanized steel sheet having a tensile strength TS of 540 MPa or more and excellent in material stability and workability (high ductility and high hole expansibility). However, I found the following.
- the positive addition of Si made it possible to improve ductility by improving the work hardening ability of ferrite, ensure strength by strengthening the solid solution of ferrite, and improve hole expansibility by relaxing the hardness difference from the second phase. Also, the use of bainitic ferrite and pearlite can alleviate the difference in hardness between soft ferrite and hard martensite, and the hole expandability can be improved.
- the present invention has been made based on the above findings, and the gist thereof is as follows.
- Component composition is mass% C: 0.04% to 0.13%, Si: 0.7% to 2.3%, Mn: 0.8% to 2.0%, P : 0.1% or less, S: 0.01% or less, Al: 0.1% or less, N: 0.008% or less, the balance consists of Fe and inevitable impurities, the steel structure has an area ratio And having a ferrite phase of 75% or more, a bainitic ferrite phase of 1.0% or more, and a pearlite phase of 1.0% or more and 10.0% or less, and the area ratio of the martensite phase is 1 It is excellent in material stability and workability characterized by satisfying a ratio of 0.0% or more and less than 5.0% and satisfying martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ⁇ 0.6 High strength hot dip galvanized steel sheet.
- a component composition it contains at least one element selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005% in mass%.
- the high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of (1) to (3).
- a component composition it contains at least one element selected from Ta: 0.001% to 0.010% and Sn: 0.002% to 0.2% by mass%.
- the high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of (1) to (4).
- a steel slab having the component composition described in any one of (1) to (6) is hot-rolled, pickled, or further cold-rolled, and then 5 ° C / ° C to a temperature range of 650 ° C or higher. heated at an average heating rate of s or more, held at a temperature range of 750 to 900 ° C. for 15 to 600 s, cooled to a temperature range of 450 to 550 ° C., and then held at a temperature range of 450 to 550 ° C. for 10 to 200 s. Then, a method for producing a high-strength hot-dip galvanized steel sheet excellent in material stability and workability, characterized by performing hot-dip galvanizing.
- the “high-strength galvanized steel sheet” is a galvanized steel sheet having a tensile strength TS of 540 MPa or more.
- the hot dip galvanized steel sheet in the present invention includes both a hot dip galvanized steel sheet that has not been subjected to an alloying treatment and an alloyed hot dip galvanized steel sheet that has been subjected to an alloying treatment.
- a high-strength hot-dip galvanized steel sheet having a tensile strength TS of 540 MPa or more and having excellent ductility and high material stability due to high ductility and high hole expansibility is obtained.
- Si was actively added for the purpose of strengthening the solid solution of ferrite and improving the work hardenability of ferrite, creating a composite structure of ferrite, bainitic ferrite, pearlite, and a small amount of martensite, and reducing the hardness difference between the different phases.
- Si was actively added for the purpose of strengthening the solid solution of ferrite and improving the work hardenability of ferrite, creating a composite structure of ferrite, bainitic ferrite, pearlite, and a small amount of martensite, and reducing the hardness difference between the different phases.
- the component composition is C: 0.04% to 0.13%, Si: 0.7% to 2.3%, Mn: 0.8% to 2.0% by mass%.
- P 0.1% or less
- S 0.01% or less
- Al 0.1% or less
- N 0.008% or less
- the ratio is 1.0% or more and less than 5.0%
- the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ⁇ 0.6 is satisfied.
- C 0.04% or more and 0.13% or less
- C is an austenite generating element and an element indispensable for strengthening steel. If the C content is less than 0.04%, it is difficult to ensure the desired strength. On the other hand, if the amount of C exceeds 0.13% and is added excessively, the welded part and the heat-affected zone are significantly hardened, and the mechanical properties of the welded part are deteriorated, so that spot weldability, arc weldability, etc. are reduced. . Therefore, C is made 0.04% or more and 0.13% or less.
- Si 0.7% or more and 2.3% or less Si is a ferrite-forming element and also an element effective for solid solution strengthening. In order to ensure good ductility by improving the work hardening ability of the ferrite phase, it is necessary to add 0.7% or more. Further, in order to ensure the desired area ratio of the bainitic ferrite phase and ensure good hole expansibility, addition of 0.7% or more is necessary. However, excessive addition of Si causes deterioration of surface properties, plating adhesion, and adhesion due to generation of red scale and the like. Therefore, Si is made 0.7% to 2.3%. Preferably, it is 1.2% or more and 1.8% or less.
- Mn 0.8% or more and 2.0% or less
- Mn is an element effective for strengthening steel.
- it is an element that stabilizes austenite, and is an element necessary for adjusting the fraction of the second phase. For this reason, it is necessary to add 0.8% or more of Mn.
- Mn is made 0.8% or more and 2.0% or less. Preferably they are 1.0% or more and 1.8% or less.
- P 0.1% or less
- P is an element effective for strengthening steel.
- P is set to 0.1% or less.
- S 0.01% or less S is an inclusion such as MnS, which causes deterioration in impact resistance and cracks along the metal flow of the weld.
- To S is set to 0.01% or less.
- Al 0.1% or less When Al exceeds 0.1%, coarse Al 2 O 3 is generated and the material deteriorates. In addition, when Al is added for deoxidation of steel, if it is less than 0.01%, many coarse oxides such as Mn and Si are dispersed in the steel and the material deteriorates. Is preferably 0.01% or more. Therefore, the Al content is 0.1% or less, preferably 0.01 to 0.1%.
- N 0.008% or less
- N is an element that causes the most deterioration of the aging resistance of the steel, and it is preferably as small as possible. If it exceeds 0.008%, the deterioration of the aging resistance becomes significant. Therefore, N is set to 0.008% or less.
- the balance is Fe and inevitable impurities.
- at least one selected from the following elements can be added as necessary.
- Cr 0.05% to 1.0%, V: 0.005% to 0.5%, Mo: 0.005% to 0.5%, Ni: 0.05% to 1.0%
- at least one selected from Cu: 0.05% or more and 1.0% or less Cr, V, and Mo has an effect of improving the balance between strength and ductility, and can be added as necessary.
- the effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more.
- excessive addition over Cr: 1.0%, V: 0.5%, and Mo: 0.5%, respectively results in an excessive fraction of the second phase, causing a significant increase in strength.
- the cost increases. Therefore, when these elements are added, the amounts are set to Cr: 1.0% or less, V: 0.5% or less, and Mo: 0.5% or less, respectively.
- Ni and Cu are effective elements for strengthening steel and can be used for strengthening steel as long as they are within the range specified in the present invention. It also has the effect of promoting internal oxidation and improving plating adhesion. In order to obtain these effects, 0.05% or more is required. On the other hand, if both Ni and Cu are added in excess of 1.0%, the workability of the steel sheet is lowered. In addition, the cost increases. Therefore, when adding Ni and Cu, the addition amount is 0.05% or more and 1.0% or less, respectively.
- B has the effect of suppressing the formation and growth of ferrite from the austenite grain boundaries, and can be added as necessary.
- the effect is obtained at 0.0003% or more.
- it exceeds 0.0050% the workability deteriorates.
- the cost increases. Therefore, when adding B, it is made into 0.0003% or more and 0.0050% or less.
- Ca 0.001% or more and 0.005% or less
- REM at least one selected from 0.001% or more and 0.005% or less
- Ta at least one selected from 0.001 to 0.010% and Sn: 0.002 to 0.2% Ta, like Ti and Nb, forms high alloy carbide and alloy carbonitride. Not only contributes to strengthening, but also partially dissolves in Nb carbide and Nb carbonitride to form a composite precipitate such as (Nb, Ta) (C, N), thereby coarsening the precipitate It is considered that there is an effect of stabilizing the contribution to strength by precipitation strengthening. Therefore, when Ta is added, the content is preferably 0.001% or more. However, if added excessively, not only the above-mentioned precipitate stabilization effect is saturated but also the alloy cost increases. Therefore, when Ta is added, its content is preferably 0.010% or less. .
- Sn can be added from the viewpoint of suppressing decarburization in the region of several tens of ⁇ m of the steel sheet surface layer caused by nitriding, oxidation, or oxidation of the steel sheet surface. By suppressing such nitriding and oxidation, the amount of martensite generated on the steel sheet surface is prevented from decreasing, and fatigue characteristics and aging resistance are improved. From the viewpoint of suppressing nitriding and oxidation, when adding Sn, its content is preferably 0.002% or more, and if it exceeds 0.2%, the toughness is reduced, so its content is reduced to 0. .2% or less is desirable.
- Sb 0.002 to 0.2%
- Sb can also be added from the viewpoint of suppressing decarburization in the region of several tens of ⁇ m of the steel sheet surface layer caused by nitridation, oxidation, or oxidation of the steel sheet surface, similarly to Sn. By suppressing such nitriding and oxidation, the amount of martensite generated on the steel sheet surface is prevented from decreasing, and fatigue characteristics and aging resistance are improved. From the viewpoint of suppressing nitriding and oxidation, when Sb is added, its content is preferably 0.002% or more, and if it exceeds 0.2%, the toughness is reduced, so the content is reduced to 0. .2% or less is desirable.
- Area ratio of ferrite phase 75% or more In order to ensure good ductility, the ferrite phase needs to have an area ratio of 75% or more.
- Area ratio of bainitic ferrite phase 1.0% or more Area ratio of bainitic ferrite phase to ensure good hole expansibility, that is, to reduce the hardness difference between soft ferrite and hard martensite 1.0% or more is necessary.
- Area ratio of pearlite phase 1.0% or more and 10.0% or less In order to ensure good hole expansibility, the area ratio of the pearlite phase is 1.0% or more. In order to secure a desired strength-ductility balance, the area ratio of the pearlite phase is set to 10.0% or less.
- reduce the amount of martensite that causes the material variation in the phase structure of the second phase and increase the amount of bainitic ferrite and pearlite that are softer than martensite. That is, it is necessary to satisfy the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ⁇ 0.6.
- carbides such as retained austenite, tempered martensite, and cementite may be generated.
- the area of the above ferrite, bainitic ferrite, pearlite, and martensite If the rate is satisfied, the object of the present invention can be achieved.
- the area ratio of ferrite, bainitic ferrite, pearlite, and martensite in the present invention is the area ratio of each phase in the observation area.
- the high-strength hot-dip galvanized steel sheet of the present invention uses a steel sheet having the above component composition and the above steel structure as a base steel sheet, and a plated film obtained by hot-dip galvanizing, or a plated film that has been subjected to alloying treatment after hot-dip galvanizing. Have.
- the high-strength hot-dip galvanized steel sheet according to the present invention is obtained by hot rolling, pickling, or further cold rolling a steel slab having a component composition suitable for the above-described component composition range, and thereafter up to a temperature range of 650 ° C. or higher.
- Heat at an average heating rate of 5 ° C./s or more hold for 15 to 600 s in a temperature range of 750 to 900 ° C., cool to a temperature range of 450 to 550 ° C., and 10 to 10 in the temperature range of 450 to 550 ° C. It is manufactured by holding for 200 s and then applying hot dip galvanizing.
- the steel having the above composition is melted by a known method, and then slab is formed through a block or continuous casting, and is hot-rolled into a hot-rolled sheet.
- hot rolling it is preferable to heat the slab to 1100 to 1300 ° C., perform hot rolling at a final finishing temperature of 850 ° C. or higher, and wind it on a steel strip at 400 to 650 ° C.
- the coiling temperature exceeds 650 ° C.
- the carbides in the hot-rolled sheet are coarsened, and such coarsened carbides cannot be melted during soaking at the time of annealing, so that the required strength may not be obtained.
- pickling treatment is performed by a known method. Or after pickling, it cold-rolls further.
- cold rolling it is not necessary to specifically limit the conditions, but it is preferable to perform cold rolling at a cold reduction rate of 30% or more. If the cold rolling reduction is low, recrystallization of ferrite is not promoted, unrecrystallized ferrite remains, and ductility and hole expansibility may decrease.
- the hot-rolled sheet pickled or cold-rolled steel sheet is subjected to the following annealing, cooled, and then hot-dip galvanized.
- anneal to hold for 15 to 600 s in the austenite single-phase region or the two-phase region of austenite and ferrite. When the annealing temperature is less than 750 ° C. and the holding time is less than 15 s, the hard cementite in the steel sheet is not sufficiently dissolved, the hole expandability is lowered, and further, the desired martensite area ratio cannot be obtained. Ductility decreases.
- the annealing temperature exceeds 900 ° C.
- the growth of austenite grains is remarkable, it becomes difficult to secure bainitic ferrite due to bainite transformation that occurs during holding after cooling, hole expansibility decreases, and the martensite area Since the ratio / (bainitic ferrite area ratio + pearlite area ratio) exceeds 0.6, good material stability cannot be obtained.
- the holding time exceeds 600 s, austenite becomes coarse, and it becomes difficult to secure a desired strength, and it may cause an increase in cost due to a large energy consumption.
- the holding temperature exceeds 550 ° C. or the holding time is less than 10 s, the bainite transformation is not promoted, the area ratio of bainitic ferrite is less than 1.0%, and the desired hole expandability cannot be obtained.
- the holding temperature is less than 450 ° C. or the holding time exceeds 200 s, most of the second phase becomes austenite and bainitic ferrite having a large amount of dissolved carbon produced by promoting bainite transformation.
- a pearlite area ratio of 0% or more cannot be obtained, and an area ratio of the hard martensite phase is 5.0% or more, and good hole expansibility and material stability cannot be obtained.
- the hot dip galvanized steel sheet in which the plated layer is not alloyed is cooled by infiltrating the steel sheet into a plating bath at a normal bath temperature and applying hot dip galvanizing. obtain.
- alloying of the plating layer In the temperature range below 500 ° C., alloying of the plating layer is not promoted, and it is difficult to obtain an galvannealed steel sheet. In the temperature range exceeding 600 ° C., most of the second phase becomes pearlite, the desired martensite area ratio cannot be obtained, and the balance between strength and ductility is lowered.
- the alloying of the plating layer can be performed without any problem within the range of the present invention that satisfies the above condition of exp [200 / (400-T)] ⁇ ln (t) in a temperature range of 500 to 600 ° C.
- the holding temperature does not need to be constant as long as it is within the above-mentioned temperature range, and even if the cooling rate changes during cooling, it may be within the specified range.
- the gist of the present invention is not impaired.
- the steel sheet may be heat-treated by any equipment.
- temper rolling of the steel sheet of the present invention for shape correction after heat treatment is also included in the scope of the present invention. In the present invention, it is assumed that the steel material is manufactured through normal steelmaking, casting, and hot rolling processes, but the manufacturing process is performed by omitting part or all of the hot rolling process by thin casting, for example. You may do it.
- 1 and 2 show No. of Steel A which is an example of the present invention in Examples described later. 15, 16, and 17 (Tables 2 and 5) and No. of Steel H as a comparative example. 18, 19, 20 (Table 2, Table 5) shows a diagram to organize TS, the relationship EL and annealing temperature (T 1). 1 and 2, it can be seen that the steel A of the present invention has a small variation in TS and EL accompanying changes in the annealing temperature, whereas the steel H of the comparative example has a large variation in TS and EL.
- 3 and 4 show the No. of steel A which is an example of the present invention in the examples described later.
- 24, 25, 26 (Table 2, Table 5) shows a diagram to organize TS, the relationship EL and the average holding temperature of the cooling after annealing (T 2).
- T 2 shows a diagram to organize TS, the relationship EL and the average holding temperature of the cooling after annealing
- the hot-rolled steel sheet (after pickling) and the cold-rolled steel sheet obtained as described above are subjected to an annealing treatment under the production conditions shown in Tables 2 to 4 by a continuous hot-dip galvanizing line, and hot-dip galvanizing treatment is performed. Furthermore, the alloying process of the plating layer was performed and the hot dip galvanized steel plate was obtained. The amount of plating was 30 to 50 g / m 2 per side. A part of a hot-dip galvanized steel sheet that was not subjected to alloying treatment after hot-dip galvanizing treatment was also produced.
- the area ratio of ferrite, bainitic ferrite, pearlite, and martensite phase with respect to the obtained hot dip galvanized steel sheet was corroded with 3% nital after polishing the plate thickness section parallel to the rolling direction of the steel sheet, and SEM Ten fields of view were observed using a (scanning electron microscope) at a magnification of 2000 times, and obtained using Image-Pro of Media Cybernetics.
- the obtained hot-dip galvanized steel sheet was subjected to tempering treatment at 200 ° C. for 2 hours, and then the structure of the plate thickness section parallel to the rolling direction of the steel sheet was described above.
- the area ratio of the tempered martensite phase obtained by the above method was defined as the area ratio of the martensite phase.
- the volume ratio of the retained austenite phase was obtained by polishing the steel plate to a 1 ⁇ 4 surface in the plate thickness direction and diffracting X-ray intensity of the 1 ⁇ 4 surface thickness. CoK ⁇ rays are used as incident X-rays, and the peaks of ⁇ 111 ⁇ , ⁇ 200 ⁇ , ⁇ 220 ⁇ , ⁇ 311 ⁇ planes of retained austenite phase and ⁇ 110 ⁇ , ⁇ 200 ⁇ , ⁇ 211 ⁇ planes of ferrite phase are used. Intensity ratios were obtained for all combinations of integrated intensities, and the average value of these ratios was taken as the volume ratio of the retained austenite phase.
- the tensile test is performed in accordance with JIS Z 2241 using a JIS No. 5 test piece obtained by taking a sample so that the tensile direction is perpendicular to the rolling direction of the steel sheet, and TS (tensile strength), EL (all Elongation) was measured.
- TS tensile strength
- EL all Elongation
- Limit hole expansion ratio ⁇ (%) ⁇ (D f ⁇ D 0 ) / D 0 ⁇ ⁇ 100
- D f hole diameter at crack initiation (mm) D 0 is the initial hole diameter (mm).
- ⁇ ⁇ 70 (%) was determined to be good.
- Each of the high-strength hot-dip galvanized steel sheets of the present invention has a TS of 540 MPa or more, an excellent hole expansibility when ⁇ is 70% or more, and a high balance of strength and ductility when TS ⁇ EL ⁇ 19000 MPa ⁇ %. It can be seen that this is a high-strength hot-dip galvanized steel sheet with excellent workability. Furthermore, it can be seen that the values of ⁇ TS and ⁇ EL are small, and the steel sheet is a high-strength hot-dip galvanized steel sheet excellent in material stability. On the other hand, in the comparative example, one or more of ductility and hole expansibility are inferior, or material stability is not preferable.
- the high-strength hot-dip galvanized steel sheet of the present invention has a tensile strength TS of 540 MPa or more, has high ductility and high hole expansibility, and is excellent in material stability.
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Abstract
Provided are a high-strength hot-dip galvanized steel sheet which has a tensile strength TS of 540 MPa or higher and excellent material stability and processability (high ductility and high hole expansibility) and a process for producing the steel sheet. The high-strength hot-dip galvanized steel sheet, which has excellent material stability and processability, is characterized by having a composition which contains, in terms of mass%, 0.04-0.13% C, 0.7-2.3% Si, 0.8-2.0% Mn, up to 0.1% P, up to 0.01% S, up to 0.1% Al, and up to 0.008% N, with the remainder comprising Fe and incidental impurities, and by having a steel structure which comprises, in terms of areal proportion, at least 75% ferrite phase, at least 1.0% bainitic ferrite phase, and 1.0-10.0% pearlite phase, has a martensite phase content of 1.0-5.0%, excluding 5.0%, in terms of areal proportion, and satisfies (areal proportion of martensite)/((areal proportion of bainitic ferrite)+(areal proportion of pearlite))≤0.6.
Description
本発明は、自動車、電気等の産業分野で使用される部材として好適な材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板およびその製造方法に関する。
The present invention relates to a high-strength hot-dip galvanized steel sheet excellent in material stability and workability suitable as a member used in industrial fields such as automobiles and electricity, and a method for producing the same.
近年、地球環境保全の見地から、自動車の燃費向上が重要な課題となっている。これに伴い、車体材料の高強度化により薄肉化を図り、車体そのものを軽量化しようとする動きが活発となってきている。
In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. Along with this, there is an active movement to reduce the thickness of the vehicle body by increasing the strength of the vehicle body material and to reduce the weight of the vehicle body itself.
しかしながら、鋼板の高強度化は延性の低下、即ち成形加工性の低下を招く。このため、高強度と高加工性を併せ持つ材料の開発が望まれているのが現状である。
However, increasing the strength of the steel sheet causes a decrease in ductility, that is, a decrease in forming processability. For this reason, the present situation is that development of a material having both high strength and high workability is desired.
また、高強度鋼板を自動車部品のような複雑な形状へ成形加工する際には、張出し部位や伸びフランジ部位で割れやネッキングの発生が大きな問題となる。そのため、割れやネッキングの発生の問題を克服できる高延性と高穴拡げ性を両立した高強度鋼板も必要とされている。
Also, when a high-strength steel sheet is formed into a complicated shape such as an automobile part, the occurrence of cracks and necking at the overhanging part and the stretched flange part becomes a big problem. Therefore, there is a need for a high-strength steel sheet that has both high ductility and high hole expansibility that can overcome the problems of cracking and necking.
さらに、鋼板の高強度化、薄肉化により形状凍結性は著しく低下する。これに対応するため、プレス成形時に、離型後の形状変化を予め予測し、形状変化量を見込んで型を設計することが広く行われているが、鋼板の引張強度(TS)が変化すると、これらを一定とした見込み量からのズレが大きくなり、形状不良が発生し、プレス成形後に一個一個形状を板金加工する等の手直しが不可欠となり、量産効率を著しく低下させる。従って、鋼板のTSのバラツキは可能な限り小さくすることが要求されている。
Furthermore, the shape freezing property is significantly reduced by increasing the strength and thinning of the steel sheet. In order to cope with this, it is widely performed to predict the shape change after mold release in advance during press forming and to design the mold in consideration of the amount of shape change, but when the tensile strength (TS) of the steel sheet changes. The deviation from the expected amount with a constant value increases, shape defects occur, reworking such as sheet metal processing one by one after press molding becomes indispensable, and mass production efficiency is significantly reduced. Therefore, it is required to make the variation in TS of the steel sheet as small as possible.
高強度鋼板の成形性向上に対しては、これまでにフェライト−マルテンサイト二相鋼(Dual−Phase鋼)や残留オーステナイトの変態誘起塑性(Transformation Induced Plasticity)を利用したTRIP鋼など、種々の複合組織型高強度溶融亜鉛めっき鋼板が開発されてきた。
To improve the formability of high-strength steel sheets, various composites such as ferritic-martensitic duplex steels (dual-phase steels) and TRIP steels utilizing transformation induced plasticity of retained austenite have been used so far. Structure-type high-strength hot-dip galvanized steel sheets have been developed.
例えば、特許文献1では、化学成分を規定し、残留オーステナイトおよびマルテンサイトの体積率、また、その製造方法を規定することにより、延性に優れた鋼板が提案されている。また、特許文献2では、化学成分を規定し、さらにその特殊な製造方法を規定することにより延性に優れた鋼板が提案されている。特許文献3では、化学成分を規定し、フェライトとベイニティックフェライトと残留オーステナイトの体積率を規定することにより、延性に優れた鋼板が提案されている。また、特許文献4では、板幅方向における伸びのバラツキが改善された高強度冷延鋼板の製造方法が提案されている。
For example, Patent Document 1 proposes a steel sheet having excellent ductility by specifying chemical components, volume ratios of retained austenite and martensite, and manufacturing methods thereof. Moreover, in patent document 2, the steel plate excellent in ductility is proposed by prescribing | regulating a chemical component and also the special manufacturing method. Patent Document 3 proposes a steel sheet having excellent ductility by defining chemical components and defining volume fractions of ferrite, bainitic ferrite and retained austenite. Patent Document 4 proposes a method for manufacturing a high-strength cold-rolled steel sheet in which variation in elongation in the sheet width direction is improved.
しかしながら、特許文献1~3では、高強度薄鋼板の延性を向上させることを主目的としているため、穴拡げ性については考慮されていない。特許文献4では、板幅方向における全伸びELのバラツキについてのみ述べており、成分組成や製造条件による材質のバラツキについては考慮されていない。そのため、高延性と高穴拡げ性を兼ね備え、かつ、材質安定性に優れた高強度溶融亜鉛めっき鋼板の開発が課題となる。
However, in Patent Documents 1 to 3, the main purpose is to improve the ductility of the high-strength thin steel sheet, and therefore the hole expandability is not considered. Patent Document 4 describes only the variation of the total elongation EL in the plate width direction, and does not consider the variation of the material due to the component composition and manufacturing conditions. Therefore, the development of a high-strength hot-dip galvanized steel sheet having both high ductility and high hole expansibility and excellent material stability becomes an issue.
本発明は、かかる事情に鑑み540MPa以上の引張強度TSを有し、かつ、材質安定性と加工性(高延性と高穴拡げ性)に優れた高強度溶融亜鉛めっき鋼板およびその製造方法を提供することを目的とする。
In view of such circumstances, the present invention provides a high-strength hot-dip galvanized steel sheet having a tensile strength TS of 540 MPa or more and excellent in material stability and workability (high ductility and high hole expansibility) and a method for producing the same. The purpose is to do.
本発明者らは、540MPa以上の引張強度TSを有し、かつ、材質安定性と加工性(高延性と高穴拡げ性)に優れた高強度溶融亜鉛めっき鋼板を得るべく鋭意検討を重ねたところ、以下のことを見出した。
The inventors have made extensive studies to obtain a high-strength hot-dip galvanized steel sheet having a tensile strength TS of 540 MPa or more and excellent in material stability and workability (high ductility and high hole expansibility). However, I found the following.
Siの積極添加により、フェライトの加工硬化能向上による延性の向上と、フェライトの固溶強化による強度確保および第二相との硬度差緩和による穴拡げ性の向上が可能となった。また、ベイニティックフェライトやパーライトの活用により、軟質なフェライトと硬質なマルテンサイトの硬度差を緩和でき、穴拡げ性の向上が可能となった。さらに、最終組織に硬質なマルテンサイトが多く存在すると軟質なフェライト相の異相界面で大きな硬度差が生じ、穴拡げ性が低下するため、最終的にマルテンサイトに変態する未変態オーステナイトをパーライト化し、フェライト、ベイニティックフェライト、パーライト、少量のマルテンサイトを有する組織を造り込むことで、高延性を維持したままで、穴拡げ性の向上が可能となり、さらに、上記各相の面積率を適正に制御することにより、材質安定性の確保が可能となった。
The positive addition of Si made it possible to improve ductility by improving the work hardening ability of ferrite, ensure strength by strengthening the solid solution of ferrite, and improve hole expansibility by relaxing the hardness difference from the second phase. Also, the use of bainitic ferrite and pearlite can alleviate the difference in hardness between soft ferrite and hard martensite, and the hole expandability can be improved. Furthermore, if there is a lot of hard martensite in the final structure, a large hardness difference occurs at the heterogeneous interface of the soft ferrite phase, and the hole expandability decreases, so the untransformed austenite that finally transforms into martensite becomes pearlite, By building a structure with ferrite, bainitic ferrite, pearlite, and a small amount of martensite, it is possible to improve hole expansibility while maintaining high ductility. By controlling, material stability can be secured.
本発明は、以上の知見に基づいてなされたものであり、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:0.04%以上0.13%以下、Si:0.7%以上2.3%以下、Mn:0.8%以上2.0%以下、P:0.1%以下、S:0.01%以下、Al:0.1%以下、N:0.008%以下を含有し、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、75%以上のフェライト相と、1.0%以上のベイニティックフェライト相と、1.0%以上10.0%以下のパーライト相を有し、さらに、マルテンサイト相の面積率が1.0%以上5.0%未満で、かつ、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6を満たすことを特徴とする材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(1) Component composition is mass% C: 0.04% to 0.13%, Si: 0.7% to 2.3%, Mn: 0.8% to 2.0%, P : 0.1% or less, S: 0.01% or less, Al: 0.1% or less, N: 0.008% or less, the balance consists of Fe and inevitable impurities, the steel structure has an area ratio And having a ferrite phase of 75% or more, a bainitic ferrite phase of 1.0% or more, and a pearlite phase of 1.0% or more and 10.0% or less, and the area ratio of the martensite phase is 1 It is excellent in material stability and workability characterized by satisfying a ratio of 0.0% or more and less than 5.0% and satisfying martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6 High strength hot dip galvanized steel sheet.
(2)さらに、成分組成として、質量%で、Cr:0.05%以上1.0%以下、V:0.005%以上0.5%以下、Mo:0.005%以上0.5%以下、Ni:0.05%以上1.0%以下、Cu:0.05%以上1.0%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ことを特徴とする(1)に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(2) Furthermore, as a component composition, Cr: 0.05% to 1.0%, V: 0.005% to 0.5%, Mo: 0.005% to 0.5% in mass% (1), characterized in that it contains at least one element selected from Ni: 0.05% to 1.0% and Cu: 0.05% to 1.0%. High-strength hot-dip galvanized steel sheet with excellent material stability and workability.
(3)さらに、成分組成として、質量%で、Ti:0.01%以上0.1%以下、Nb:0.01%以上0.1%以下、B:0.0003%以上0.0050%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ことを特徴とする(1)または(2)に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(3) Further, as a component composition, Ti: 0.01% to 0.1%, Nb: 0.01% to 0.1%, B: 0.0003% to 0.0050% in mass% The high-strength hot-dip galvanized steel sheet having excellent material stability and workability according to (1) or (2), comprising at least one element selected from the following.
(4)さらに、成分組成として、質量%で、Ca:0.001%以上0.005%以下、REM:0.001%以上0.005%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ことを特徴とする(1)~(3)のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(4) Furthermore, as a component composition, it contains at least one element selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005% in mass%. The high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of (1) to (3).
(5)さらに、成分組成として、質量%で、Ta:0.001%以上0.010%以下、Sn:0.002%以上0.2%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ことを特徴とする(1)~(4)のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(5) Further, as a component composition, it contains at least one element selected from Ta: 0.001% to 0.010% and Sn: 0.002% to 0.2% by mass%. The high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of (1) to (4).
(6)さらに、成分組成として、質量%で、Sb:0.002%以上0.2%以下を含有することを特徴とする(1)~(5)のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。
(6) The material stability according to any one of (1) to (5), further comprising, as a component composition, Sb: 0.002% to 0.2% by mass% High-strength hot-dip galvanized steel sheet with excellent workability.
(7) (1)~(6)のいずれかに記載の成分組成を有する鋼スラブを、熱間圧延、酸洗し、またはさらに冷間圧延し、その後650℃以上の温度域まで5℃/s以上の平均加熱速度で加熱し、750~900℃の温度域で15~600s保持し、450~550℃の温度域に冷却した後、該450~550℃の温度域で10~200s保持し、次いで、溶融亜鉛めっきを施すことを特徴とする材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
(7) A steel slab having the component composition described in any one of (1) to (6) is hot-rolled, pickled, or further cold-rolled, and then 5 ° C / ° C to a temperature range of 650 ° C or higher. heated at an average heating rate of s or more, held at a temperature range of 750 to 900 ° C. for 15 to 600 s, cooled to a temperature range of 450 to 550 ° C., and then held at a temperature range of 450 to 550 ° C. for 10 to 200 s. Then, a method for producing a high-strength hot-dip galvanized steel sheet excellent in material stability and workability, characterized by performing hot-dip galvanizing.
(8)溶融亜鉛めっきを施した後、500~600℃の温度域において下式を満たす条件で亜鉛めっきの合金化処理を施すことを特徴とする(7)に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 (8) The material stability and workability as described in (7), characterized in that after hot dip galvanization, galvanization is alloyed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula: For producing high-strength hot-dip galvanized steel sheets with excellent resistance.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 (8) The material stability and workability as described in (7), characterized in that after hot dip galvanization, galvanization is alloyed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula: For producing high-strength hot-dip galvanized steel sheets with excellent resistance.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
なお、本明細書において、鋼の成分を示す%は、すべて質量%である。また、本発明において、「高強度溶融亜鉛めっき鋼板」とは、引張強度TSが540MPa以上である溶融亜鉛めっき鋼板である。
In addition, in this specification, all% which shows the component of steel is the mass%. In the present invention, the “high-strength galvanized steel sheet” is a galvanized steel sheet having a tensile strength TS of 540 MPa or more.
また、本発明においては、合金化処理を施す、施さないにかかわらず、溶融亜鉛めっきによって鋼板上に亜鉛をめっきした鋼板を総称して溶融亜鉛めっき鋼板と呼称する。すなわち、本発明における溶融亜鉛めっき鋼板とは、合金化処理を施してない溶融亜鉛めっき鋼板、合金化処理を施した合金化溶融亜鉛めっき鋼板の両方を含むものである。
In the present invention, regardless of whether or not the alloying treatment is performed, a steel plate in which zinc is plated on the steel plate by hot dip galvanizing is generically called a hot dip galvanized steel plate. That is, the hot dip galvanized steel sheet in the present invention includes both a hot dip galvanized steel sheet that has not been subjected to an alloying treatment and an alloyed hot dip galvanized steel sheet that has been subjected to an alloying treatment.
本発明によれば、540MPa以上の引張強度TSを有し、かつ、高延性と高穴拡げ性であることから加工性に優れ、さらに材質安定性に優れた高強度溶融亜鉛めっき鋼板が得られる。本発明の高強度溶融亜鉛めっき鋼板を、例えば、自動車構造部材に適用することにより車体軽量化による燃費改善を図ることができ、産業上の利用価値は非常に大きい。
According to the present invention, a high-strength hot-dip galvanized steel sheet having a tensile strength TS of 540 MPa or more and having excellent ductility and high material stability due to high ductility and high hole expansibility is obtained. . By applying the high-strength hot-dip galvanized steel sheet of the present invention to, for example, an automobile structural member, fuel efficiency can be improved by reducing the weight of the vehicle body, and the industrial utility value is very large.
以下に、本発明の詳細を説明する。
Details of the present invention will be described below.
一般に、軟質なフェライトと硬質なマルテンサイトとの二相構造では、延性の確保は可能なものの、フェライトとマルテンサイトの硬度差が大きいために、十分な穴拡げ性が得られないことが知られている。そこで、本発明者は、さらにベイニティックフェライトとパーライトの活用について検討し、フェライトとベイニティックフェライトとパーライトとマルテンサイトを有する(一部残留オーステナイトを有するものを含む)複合組織での特性向上の可能性に着目して詳細に検討を行った。
In general, in a two-phase structure of soft ferrite and hard martensite, it is possible to ensure ductility, but it is known that sufficient hole expansibility cannot be obtained due to the large hardness difference between ferrite and martensite. ing. Therefore, the present inventor further examined the utilization of bainitic ferrite and pearlite, and improved the characteristics in the composite structure including ferrite, bainitic ferrite, pearlite, and martensite (including those having some retained austenite). We examined in detail the possibility of this.
その結果、フェライトの固溶強化とフェライトの加工硬化能向上を目的にSiを積極添加し、フェライトとベイニティックフェライトとパーライトと少量のマルテンサイトの複合組織を造り込み、異相間の硬度差を低減させ、さらにその複合組織の面積分率を適正化することにより、高延性と高穴拡げ性の両立、及び、材質安定性の確保が可能となった。
As a result, Si was actively added for the purpose of strengthening the solid solution of ferrite and improving the work hardenability of ferrite, creating a composite structure of ferrite, bainitic ferrite, pearlite, and a small amount of martensite, and reducing the hardness difference between the different phases. By reducing and further optimizing the area fraction of the composite structure, it became possible to achieve both high ductility and high hole expansibility and to ensure material stability.
以上が本発明を完成するに至った技術的特徴である。そして、本発明は、成分組成は、質量%でC:0.04%以上0.13%以下、Si:0.7%以上2.3%以下、Mn:0.8%以上2.0%以下、P:0.1%以下、S:0.01%以下、Al:0.1%以下、N:0.008%以下を含有し、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、75%以上のフェライト相と、1.0%以上のベイニティックフェライト相と、1.0%以上10.0%以下のパーライト相を有し、さらに、マルテンサイト相の面積率が1.0%以上5.0%未満で、かつ、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6を満たすことを特徴とする。
These are the technical features that led to the completion of the present invention. In the present invention, the component composition is C: 0.04% to 0.13%, Si: 0.7% to 2.3%, Mn: 0.8% to 2.0% by mass%. Hereinafter, P: 0.1% or less, S: 0.01% or less, Al: 0.1% or less, N: 0.008% or less, with the balance being Fe and inevitable impurities, And having an area ratio of 75% or more of ferrite phase, 1.0% or more of bainitic ferrite phase, 1.0% or more and 10.0% or less of pearlite phase, and the area of martensite phase The ratio is 1.0% or more and less than 5.0%, and the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6 is satisfied.
(1)まず、成分組成について説明する。
(1) First, the component composition will be described.
C:0.04%以上0.13%以下
Cはオーステナイト生成元素であり、鋼の強化に不可欠な元素である。C量が0.04%未満では、所望の強度確保が難しい。一方、C量が0.13%を超えて過剰に添加すると、溶接部および熱影響部の硬化が著しく、溶接部の機械的特性が劣化するため、スポット溶接性、アーク溶接性等が低下する。よって、Cは0.04%以上0.13%以下とする。 C: 0.04% or more and 0.13% or less C is an austenite generating element and an element indispensable for strengthening steel. If the C content is less than 0.04%, it is difficult to ensure the desired strength. On the other hand, if the amount of C exceeds 0.13% and is added excessively, the welded part and the heat-affected zone are significantly hardened, and the mechanical properties of the welded part are deteriorated, so that spot weldability, arc weldability, etc. are reduced. . Therefore, C is made 0.04% or more and 0.13% or less.
Cはオーステナイト生成元素であり、鋼の強化に不可欠な元素である。C量が0.04%未満では、所望の強度確保が難しい。一方、C量が0.13%を超えて過剰に添加すると、溶接部および熱影響部の硬化が著しく、溶接部の機械的特性が劣化するため、スポット溶接性、アーク溶接性等が低下する。よって、Cは0.04%以上0.13%以下とする。 C: 0.04% or more and 0.13% or less C is an austenite generating element and an element indispensable for strengthening steel. If the C content is less than 0.04%, it is difficult to ensure the desired strength. On the other hand, if the amount of C exceeds 0.13% and is added excessively, the welded part and the heat-affected zone are significantly hardened, and the mechanical properties of the welded part are deteriorated, so that spot weldability, arc weldability, etc. are reduced. . Therefore, C is made 0.04% or more and 0.13% or less.
Si:0.7%以上2.3%以下
Siはフェライト生成元素であり、また、固溶強化に有効な元素でもある。そして、フェライト相の加工硬化能向上による良好な延性確保のためには0.7%以上の添加が必要である。さらに、所望のベイニティックフェライト相の面積率を確保し、良好な穴拡げ性を確保するためには0.7%以上の添加が必要である。しかしながら、Siの過剰な添加は、赤スケール等の発生により表面性状の劣化や、めっき付着・密着性の劣化を引き起こす。よって、Siは0.7%以上2.3%以下とする。好ましくは、1.2%以上1.8%以下である。 Si: 0.7% or more and 2.3% or less Si is a ferrite-forming element and also an element effective for solid solution strengthening. In order to ensure good ductility by improving the work hardening ability of the ferrite phase, it is necessary to add 0.7% or more. Further, in order to ensure the desired area ratio of the bainitic ferrite phase and ensure good hole expansibility, addition of 0.7% or more is necessary. However, excessive addition of Si causes deterioration of surface properties, plating adhesion, and adhesion due to generation of red scale and the like. Therefore, Si is made 0.7% to 2.3%. Preferably, it is 1.2% or more and 1.8% or less.
Siはフェライト生成元素であり、また、固溶強化に有効な元素でもある。そして、フェライト相の加工硬化能向上による良好な延性確保のためには0.7%以上の添加が必要である。さらに、所望のベイニティックフェライト相の面積率を確保し、良好な穴拡げ性を確保するためには0.7%以上の添加が必要である。しかしながら、Siの過剰な添加は、赤スケール等の発生により表面性状の劣化や、めっき付着・密着性の劣化を引き起こす。よって、Siは0.7%以上2.3%以下とする。好ましくは、1.2%以上1.8%以下である。 Si: 0.7% or more and 2.3% or less Si is a ferrite-forming element and also an element effective for solid solution strengthening. In order to ensure good ductility by improving the work hardening ability of the ferrite phase, it is necessary to add 0.7% or more. Further, in order to ensure the desired area ratio of the bainitic ferrite phase and ensure good hole expansibility, addition of 0.7% or more is necessary. However, excessive addition of Si causes deterioration of surface properties, plating adhesion, and adhesion due to generation of red scale and the like. Therefore, Si is made 0.7% to 2.3%. Preferably, it is 1.2% or more and 1.8% or less.
Mn:0.8%以上2.0%以下
Mnは、鋼の強化に有効な元素である。また、オーステナイトを安定化させる元素であり、第二相の分率調整に必要な元素である。このため、Mnは0.8%以上の添加が必要である。一方、2.0%を超えて過剰に添加すると、第二相中のマルテンサイト面積率が増加し、材質安定性の確保が困難となる。また、近年Mnの合金コストが高騰しているため、コストアップの要因にも繋がる。従って、Mnは0.8%以上2.0%以下とする。好ましくは1.0%以上1.8%以下である。 Mn: 0.8% or more and 2.0% or less Mn is an element effective for strengthening steel. In addition, it is an element that stabilizes austenite, and is an element necessary for adjusting the fraction of the second phase. For this reason, it is necessary to add 0.8% or more of Mn. On the other hand, when it exceeds 2.0% and it adds excessively, the martensite area rate in a 2nd phase will increase, and ensuring of material stability will become difficult. Moreover, since the alloy cost of Mn has soared in recent years, it also leads to a cost increase factor. Therefore, Mn is made 0.8% or more and 2.0% or less. Preferably they are 1.0% or more and 1.8% or less.
Mnは、鋼の強化に有効な元素である。また、オーステナイトを安定化させる元素であり、第二相の分率調整に必要な元素である。このため、Mnは0.8%以上の添加が必要である。一方、2.0%を超えて過剰に添加すると、第二相中のマルテンサイト面積率が増加し、材質安定性の確保が困難となる。また、近年Mnの合金コストが高騰しているため、コストアップの要因にも繋がる。従って、Mnは0.8%以上2.0%以下とする。好ましくは1.0%以上1.8%以下である。 Mn: 0.8% or more and 2.0% or less Mn is an element effective for strengthening steel. In addition, it is an element that stabilizes austenite, and is an element necessary for adjusting the fraction of the second phase. For this reason, it is necessary to add 0.8% or more of Mn. On the other hand, when it exceeds 2.0% and it adds excessively, the martensite area rate in a 2nd phase will increase, and ensuring of material stability will become difficult. Moreover, since the alloy cost of Mn has soared in recent years, it also leads to a cost increase factor. Therefore, Mn is made 0.8% or more and 2.0% or less. Preferably they are 1.0% or more and 1.8% or less.
P:0.1%以下
Pは、鋼の強化に有効な元素であるが、0.1%を超えて過剰に添加すると、粒界偏析により脆化を引き起こし、耐衝撃性を劣化させる。また0.1%を超えると合金化速度を大幅に遅延させる。従って、Pは0.1%以下とする。 P: 0.1% or less P is an element effective for strengthening steel. However, when P is added excessively in excess of 0.1%, embrittlement occurs due to segregation at the grain boundaries and impact resistance is deteriorated. If it exceeds 0.1%, the alloying speed is significantly delayed. Therefore, P is set to 0.1% or less.
Pは、鋼の強化に有効な元素であるが、0.1%を超えて過剰に添加すると、粒界偏析により脆化を引き起こし、耐衝撃性を劣化させる。また0.1%を超えると合金化速度を大幅に遅延させる。従って、Pは0.1%以下とする。 P: 0.1% or less P is an element effective for strengthening steel. However, when P is added excessively in excess of 0.1%, embrittlement occurs due to segregation at the grain boundaries and impact resistance is deteriorated. If it exceeds 0.1%, the alloying speed is significantly delayed. Therefore, P is set to 0.1% or less.
S:0.01%以下
Sは、MnSなどの介在物となって、耐衝撃性の劣化や溶接部のメタルフローに沿った割れの原因となるので極力低い方がよいが、製造コストの面からSは0.01%以下とする。 S: 0.01% or less S is an inclusion such as MnS, which causes deterioration in impact resistance and cracks along the metal flow of the weld. To S is set to 0.01% or less.
Sは、MnSなどの介在物となって、耐衝撃性の劣化や溶接部のメタルフローに沿った割れの原因となるので極力低い方がよいが、製造コストの面からSは0.01%以下とする。 S: 0.01% or less S is an inclusion such as MnS, which causes deterioration in impact resistance and cracks along the metal flow of the weld. To S is set to 0.01% or less.
Al:0.1%以下
Alは、0.1%を超えると、粗大なAl2O3が生成し、材質が劣化する。また、Alは鋼の脱酸のために添加される場合、0.01%未満ではMnやSiなどの粗大な酸化物が鋼中に多数分散して材質が劣化することになるため、添加量を0.01%以上とするのが好ましい。よって、Al量は0.1%以下とし、好ましくは、0.01~0.1%とする。 Al: 0.1% or less When Al exceeds 0.1%, coarse Al 2 O 3 is generated and the material deteriorates. In addition, when Al is added for deoxidation of steel, if it is less than 0.01%, many coarse oxides such as Mn and Si are dispersed in the steel and the material deteriorates. Is preferably 0.01% or more. Therefore, the Al content is 0.1% or less, preferably 0.01 to 0.1%.
Alは、0.1%を超えると、粗大なAl2O3が生成し、材質が劣化する。また、Alは鋼の脱酸のために添加される場合、0.01%未満ではMnやSiなどの粗大な酸化物が鋼中に多数分散して材質が劣化することになるため、添加量を0.01%以上とするのが好ましい。よって、Al量は0.1%以下とし、好ましくは、0.01~0.1%とする。 Al: 0.1% or less When Al exceeds 0.1%, coarse Al 2 O 3 is generated and the material deteriorates. In addition, when Al is added for deoxidation of steel, if it is less than 0.01%, many coarse oxides such as Mn and Si are dispersed in the steel and the material deteriorates. Is preferably 0.01% or more. Therefore, the Al content is 0.1% or less, preferably 0.01 to 0.1%.
N:0.008%以下
Nは、鋼の耐時効性を最も大きく劣化させる元素であり、少ないほど好ましく、0.008%を超えると耐時効性の劣化が顕著となる。従って、Nは0.008%以下とする。 N: 0.008% or less N is an element that causes the most deterioration of the aging resistance of the steel, and it is preferably as small as possible. If it exceeds 0.008%, the deterioration of the aging resistance becomes significant. Therefore, N is set to 0.008% or less.
Nは、鋼の耐時効性を最も大きく劣化させる元素であり、少ないほど好ましく、0.008%を超えると耐時効性の劣化が顕著となる。従って、Nは0.008%以下とする。 N: 0.008% or less N is an element that causes the most deterioration of the aging resistance of the steel, and it is preferably as small as possible. If it exceeds 0.008%, the deterioration of the aging resistance becomes significant. Therefore, N is set to 0.008% or less.
残部はFeおよび不可避的不純物である。ただし、これらの元素に加えて、以下の元素のうちから選ばれる少なくとも1種を必要に応じて添加することができる。
The balance is Fe and inevitable impurities. However, in addition to these elements, at least one selected from the following elements can be added as necessary.
Cr:0.05%以上1.0%以下、V:0.005%以上0.5%以下、Mo:0.005%以上0.5%以下、Ni:0.05%以上1.0%以下、Cu:0.05%以上1.0%以下のうちから選ばれる少なくとも1種
Cr、V、Moは強度と延性のバランスを向上させる作用を有するので必要に応じて添加することができる。その効果は、Cr:0.05%以上、V:0.005%以上、Mo:0.005%以上で得られる。しかしながら、それぞれCr:1.0%、V:0.5%、Mo:0.5%を超えて過剰に添加すると、第二相の分率が過大となり著しい強度上昇等の懸念が生じる。また、コストアップの要因にもなる。従って、これらの元素を添加する場合には、その量をそれぞれCr:1.0%以下、V:0.5%以下、Mo:0.5%以下とする。 Cr: 0.05% to 1.0%, V: 0.005% to 0.5%, Mo: 0.005% to 0.5%, Ni: 0.05% to 1.0% Hereinafter, at least one selected from Cu: 0.05% or more and 1.0% or less Cr, V, and Mo has an effect of improving the balance between strength and ductility, and can be added as necessary. The effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more. However, excessive addition over Cr: 1.0%, V: 0.5%, and Mo: 0.5%, respectively, results in an excessive fraction of the second phase, causing a significant increase in strength. In addition, the cost increases. Therefore, when these elements are added, the amounts are set to Cr: 1.0% or less, V: 0.5% or less, and Mo: 0.5% or less, respectively.
Cr、V、Moは強度と延性のバランスを向上させる作用を有するので必要に応じて添加することができる。その効果は、Cr:0.05%以上、V:0.005%以上、Mo:0.005%以上で得られる。しかしながら、それぞれCr:1.0%、V:0.5%、Mo:0.5%を超えて過剰に添加すると、第二相の分率が過大となり著しい強度上昇等の懸念が生じる。また、コストアップの要因にもなる。従って、これらの元素を添加する場合には、その量をそれぞれCr:1.0%以下、V:0.5%以下、Mo:0.5%以下とする。 Cr: 0.05% to 1.0%, V: 0.005% to 0.5%, Mo: 0.005% to 0.5%, Ni: 0.05% to 1.0% Hereinafter, at least one selected from Cu: 0.05% or more and 1.0% or less Cr, V, and Mo has an effect of improving the balance between strength and ductility, and can be added as necessary. The effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more. However, excessive addition over Cr: 1.0%, V: 0.5%, and Mo: 0.5%, respectively, results in an excessive fraction of the second phase, causing a significant increase in strength. In addition, the cost increases. Therefore, when these elements are added, the amounts are set to Cr: 1.0% or less, V: 0.5% or less, and Mo: 0.5% or less, respectively.
Ni、Cuは鋼の強化に有効な元素であり、本発明で規定した範囲内であれば鋼の強化に使用して差し支えない。また内部酸化を促進してめっき密着性を向上させる作用がある。これらの効果を得るためには,それぞれ0.05%以上必要である。一方、Ni、Cuとも1.0%を超えて添加すると、鋼板の加工性を低下させる。また、コストアップの要因にもなる。よって、Ni、Cuを添加する場合に、その添加量はそれぞれ0.05%以上1.0%以下とする。
Ni and Cu are effective elements for strengthening steel and can be used for strengthening steel as long as they are within the range specified in the present invention. It also has the effect of promoting internal oxidation and improving plating adhesion. In order to obtain these effects, 0.05% or more is required. On the other hand, if both Ni and Cu are added in excess of 1.0%, the workability of the steel sheet is lowered. In addition, the cost increases. Therefore, when adding Ni and Cu, the addition amount is 0.05% or more and 1.0% or less, respectively.
Ti:0.01%以上0.1%以下、Nb:0.01%以上0.1%以下、B:0.0003%以上0.0050%以下のうちから選ばれる少なくとも1種
Ti、Nbは鋼の析出強化に有効で、その効果はそれぞれ0.01%以上で得られ、本発明で規定した範囲内であれば鋼の強化に使用して差し支えない。しかし、それぞれが0.1%を超えると加工性および形状凍結性が低下する。また、コストアップの要因にもなる。従って、Ti、Nbを添加する場合には,その添加量をTiは0.01%以上0.1%以下、Nbは0.01%以上0.1%以下とする。 At least one selected from Ti: 0.01% to 0.1%, Nb: 0.01% to 0.1%, B: 0.0003% to 0.0050% Ti, Nb It is effective for precipitation strengthening of steel, and the effect can be obtained at 0.01% or more. If it is within the range specified by the present invention, it can be used for strengthening steel. However, when each exceeds 0.1%, workability and shape freezing property will fall. In addition, the cost increases. Therefore, when adding Ti and Nb, the addition amount is set to 0.01% to 0.1% for Ti and 0.01% to 0.1% for Nb.
Ti、Nbは鋼の析出強化に有効で、その効果はそれぞれ0.01%以上で得られ、本発明で規定した範囲内であれば鋼の強化に使用して差し支えない。しかし、それぞれが0.1%を超えると加工性および形状凍結性が低下する。また、コストアップの要因にもなる。従って、Ti、Nbを添加する場合には,その添加量をTiは0.01%以上0.1%以下、Nbは0.01%以上0.1%以下とする。 At least one selected from Ti: 0.01% to 0.1%, Nb: 0.01% to 0.1%, B: 0.0003% to 0.0050% Ti, Nb It is effective for precipitation strengthening of steel, and the effect can be obtained at 0.01% or more. If it is within the range specified by the present invention, it can be used for strengthening steel. However, when each exceeds 0.1%, workability and shape freezing property will fall. In addition, the cost increases. Therefore, when adding Ti and Nb, the addition amount is set to 0.01% to 0.1% for Ti and 0.01% to 0.1% for Nb.
Bはオーステナイト粒界からのフェライトの生成・成長を抑制する作用を有するので必要に応じて添加することができる。その効果は,0.0003%以上で得られる。しかし、0.0050%を超えると加工性が低下する。また、コストアップの要因にもなる。従って、Bを添加する場合は0.0003%以上0.0050%以下とする。
B has the effect of suppressing the formation and growth of ferrite from the austenite grain boundaries, and can be added as necessary. The effect is obtained at 0.0003% or more. However, if it exceeds 0.0050%, the workability deteriorates. In addition, the cost increases. Therefore, when adding B, it is made into 0.0003% or more and 0.0050% or less.
Ca:0.001%以上0.005%以下、REM:0.001%以上0.005%以下のうちから選ばれる少なくとも1種
CaおよびREMは、硫化物の形状を球状化し穴拡げ性への硫化物の悪影響を改善するために有効な元素である。この効果を得るためには、それぞれ0.001%以上必要である。しかしながら、過剰な添加は,介在物等の増加を引き起こし表面および内部欠陥などを引き起こす。したがって、Ca、REMを添加する場合は、その添加量はそれぞれ0.001%以上0.005%以下とする。 Ca: 0.001% or more and 0.005% or less, REM: at least one selected from 0.001% or more and 0.005% or less Ca and REM spheroidize the shape of the sulfide to improve the hole expandability It is an effective element for improving the adverse effects of sulfides. In order to obtain this effect, 0.001% or more is required for each. However, excessive addition causes an increase in inclusions and causes surface and internal defects. Therefore, when Ca and REM are added, the addition amounts are 0.001% or more and 0.005% or less, respectively.
CaおよびREMは、硫化物の形状を球状化し穴拡げ性への硫化物の悪影響を改善するために有効な元素である。この効果を得るためには、それぞれ0.001%以上必要である。しかしながら、過剰な添加は,介在物等の増加を引き起こし表面および内部欠陥などを引き起こす。したがって、Ca、REMを添加する場合は、その添加量はそれぞれ0.001%以上0.005%以下とする。 Ca: 0.001% or more and 0.005% or less, REM: at least one selected from 0.001% or more and 0.005% or less Ca and REM spheroidize the shape of the sulfide to improve the hole expandability It is an effective element for improving the adverse effects of sulfides. In order to obtain this effect, 0.001% or more is required for each. However, excessive addition causes an increase in inclusions and causes surface and internal defects. Therefore, when Ca and REM are added, the addition amounts are 0.001% or more and 0.005% or less, respectively.
Ta:0.001~0.010%、Sn:0.002~0.2%のうちから選ばれる少なくとも1種
Taは、TiやNbと同様、合金炭化物や合金炭窒化物を形成して高強度化に寄与するのみならず、Nb炭化物やNb炭窒化物に一部固溶し、(Nb,Ta)(C,N)のような複合析出物を形成することで、析出物の粗大化を著しく抑制して、析出強化による強度への寄与を安定化させる効果があると考えられる。そのため、Taを添加する場合は、その含有量を0.001%以上とすることが望ましい。しかし、過剰に添加した場合、上記の析出物安定化効果が飽和するのみならず、合金コストが上昇するため、Taを添加する場合は、その含有量を0.010%以下とすることが望ましい。 Ta: at least one selected from 0.001 to 0.010% and Sn: 0.002 to 0.2% Ta, like Ti and Nb, forms high alloy carbide and alloy carbonitride. Not only contributes to strengthening, but also partially dissolves in Nb carbide and Nb carbonitride to form a composite precipitate such as (Nb, Ta) (C, N), thereby coarsening the precipitate It is considered that there is an effect of stabilizing the contribution to strength by precipitation strengthening. Therefore, when Ta is added, the content is preferably 0.001% or more. However, if added excessively, not only the above-mentioned precipitate stabilization effect is saturated but also the alloy cost increases. Therefore, when Ta is added, its content is preferably 0.010% or less. .
Taは、TiやNbと同様、合金炭化物や合金炭窒化物を形成して高強度化に寄与するのみならず、Nb炭化物やNb炭窒化物に一部固溶し、(Nb,Ta)(C,N)のような複合析出物を形成することで、析出物の粗大化を著しく抑制して、析出強化による強度への寄与を安定化させる効果があると考えられる。そのため、Taを添加する場合は、その含有量を0.001%以上とすることが望ましい。しかし、過剰に添加した場合、上記の析出物安定化効果が飽和するのみならず、合金コストが上昇するため、Taを添加する場合は、その含有量を0.010%以下とすることが望ましい。 Ta: at least one selected from 0.001 to 0.010% and Sn: 0.002 to 0.2% Ta, like Ti and Nb, forms high alloy carbide and alloy carbonitride. Not only contributes to strengthening, but also partially dissolves in Nb carbide and Nb carbonitride to form a composite precipitate such as (Nb, Ta) (C, N), thereby coarsening the precipitate It is considered that there is an effect of stabilizing the contribution to strength by precipitation strengthening. Therefore, when Ta is added, the content is preferably 0.001% or more. However, if added excessively, not only the above-mentioned precipitate stabilization effect is saturated but also the alloy cost increases. Therefore, when Ta is added, its content is preferably 0.010% or less. .
Snは、鋼板表面の窒化、酸化、あるいは酸化により生じる鋼板表層の数10μm領域の脱炭を抑制する観点から添加することができる。このような窒化や酸化を抑制することで鋼板表面においてマルテンサイトの生成量が減少するのを防止し、疲労特性や耐時効性を改善させる。窒化や酸化を抑制する観点から、Snを添加する場合は、その含有量は0.002%以上とすることが望ましく、0.2%を超えると靭性の低下を招くため、その含有量を0.2%以下とすることが望ましい。
Sn can be added from the viewpoint of suppressing decarburization in the region of several tens of μm of the steel sheet surface layer caused by nitriding, oxidation, or oxidation of the steel sheet surface. By suppressing such nitriding and oxidation, the amount of martensite generated on the steel sheet surface is prevented from decreasing, and fatigue characteristics and aging resistance are improved. From the viewpoint of suppressing nitriding and oxidation, when adding Sn, its content is preferably 0.002% or more, and if it exceeds 0.2%, the toughness is reduced, so its content is reduced to 0. .2% or less is desirable.
Sb:0.002~0.2%
SbもSnと同様に鋼板表面の窒化、酸化、あるいは酸化により生じる鋼板表層の数10μm領域の脱炭を抑制する観点から添加することができる。このような窒化や酸化を抑制することで鋼板表面においてマルテンサイトの生成量が減少するのを防止し、疲労特性や耐時効性を改善させる。窒化や酸化を抑制する観点から、Sbを添加する場合は、その含有量は0.002%以上とすることが望ましく、0.2%を超えると靭性の低下を招くため、その含有量を0.2%以下とすることが望ましい。 Sb: 0.002 to 0.2%
Sb can also be added from the viewpoint of suppressing decarburization in the region of several tens of μm of the steel sheet surface layer caused by nitridation, oxidation, or oxidation of the steel sheet surface, similarly to Sn. By suppressing such nitriding and oxidation, the amount of martensite generated on the steel sheet surface is prevented from decreasing, and fatigue characteristics and aging resistance are improved. From the viewpoint of suppressing nitriding and oxidation, when Sb is added, its content is preferably 0.002% or more, and if it exceeds 0.2%, the toughness is reduced, so the content is reduced to 0. .2% or less is desirable.
SbもSnと同様に鋼板表面の窒化、酸化、あるいは酸化により生じる鋼板表層の数10μm領域の脱炭を抑制する観点から添加することができる。このような窒化や酸化を抑制することで鋼板表面においてマルテンサイトの生成量が減少するのを防止し、疲労特性や耐時効性を改善させる。窒化や酸化を抑制する観点から、Sbを添加する場合は、その含有量は0.002%以上とすることが望ましく、0.2%を超えると靭性の低下を招くため、その含有量を0.2%以下とすることが望ましい。 Sb: 0.002 to 0.2%
Sb can also be added from the viewpoint of suppressing decarburization in the region of several tens of μm of the steel sheet surface layer caused by nitridation, oxidation, or oxidation of the steel sheet surface, similarly to Sn. By suppressing such nitriding and oxidation, the amount of martensite generated on the steel sheet surface is prevented from decreasing, and fatigue characteristics and aging resistance are improved. From the viewpoint of suppressing nitriding and oxidation, when Sb is added, its content is preferably 0.002% or more, and if it exceeds 0.2%, the toughness is reduced, so the content is reduced to 0. .2% or less is desirable.
(2)次に鋼組織について説明する。
(2) Next, the steel structure will be explained.
フェライト相の面積率:75%以上
良好な延性を確保するためには、フェライト相は面積率で75%以上必要である。 Area ratio of ferrite phase: 75% or more In order to ensure good ductility, the ferrite phase needs to have an area ratio of 75% or more.
良好な延性を確保するためには、フェライト相は面積率で75%以上必要である。 Area ratio of ferrite phase: 75% or more In order to ensure good ductility, the ferrite phase needs to have an area ratio of 75% or more.
ベイニティックフェライト相の面積率:1.0%以上
良好な穴拡げ性の確保のため、即ち軟質なフェライトと硬質なマルテンサイトの硬度差を緩和させるために、ベイニティックフェライト相の面積率は1.0%以上必要である。 Area ratio of bainitic ferrite phase: 1.0% or more Area ratio of bainitic ferrite phase to ensure good hole expansibility, that is, to reduce the hardness difference between soft ferrite and hard martensite 1.0% or more is necessary.
良好な穴拡げ性の確保のため、即ち軟質なフェライトと硬質なマルテンサイトの硬度差を緩和させるために、ベイニティックフェライト相の面積率は1.0%以上必要である。 Area ratio of bainitic ferrite phase: 1.0% or more Area ratio of bainitic ferrite phase to ensure good hole expansibility, that is, to reduce the hardness difference between soft ferrite and hard martensite 1.0% or more is necessary.
パーライト相の面積率:1.0%以上10.0%以下
良好な穴拡げ性の確保のため、パーライト相の面積率は1.0%以上とする。所望の強度−延性バランスを確保するため、パーライト相の面積率を10.0%以下とする。 Area ratio of pearlite phase: 1.0% or more and 10.0% or less In order to ensure good hole expansibility, the area ratio of the pearlite phase is 1.0% or more. In order to secure a desired strength-ductility balance, the area ratio of the pearlite phase is set to 10.0% or less.
良好な穴拡げ性の確保のため、パーライト相の面積率は1.0%以上とする。所望の強度−延性バランスを確保するため、パーライト相の面積率を10.0%以下とする。 Area ratio of pearlite phase: 1.0% or more and 10.0% or less In order to ensure good hole expansibility, the area ratio of the pearlite phase is 1.0% or more. In order to secure a desired strength-ductility balance, the area ratio of the pearlite phase is set to 10.0% or less.
マルテンサイト相の面積率:1.0%以上5.0%未満
所望の強度−延性バランスを確保するため、マルテンサイト相の面積率は1.0%以上とする。良好な材質安定性を確保するために、引張特性(TS、EL)に大きく影響を及ぼすマルテンサイト相の面積率は5.0%未満である必要がある。 Area ratio of martensite phase: 1.0% or more and less than 5.0% In order to secure a desired strength-ductility balance, the area ratio of the martensite phase is 1.0% or more. In order to ensure good material stability, the area ratio of the martensite phase that greatly affects the tensile properties (TS, EL) needs to be less than 5.0%.
所望の強度−延性バランスを確保するため、マルテンサイト相の面積率は1.0%以上とする。良好な材質安定性を確保するために、引張特性(TS、EL)に大きく影響を及ぼすマルテンサイト相の面積率は5.0%未満である必要がある。 Area ratio of martensite phase: 1.0% or more and less than 5.0% In order to secure a desired strength-ductility balance, the area ratio of the martensite phase is 1.0% or more. In order to ensure good material stability, the area ratio of the martensite phase that greatly affects the tensile properties (TS, EL) needs to be less than 5.0%.
マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6
良好な材質安定性を確保するために、第二相の相構成を、材質バラツキの要因となるマルテンサイトの量を低減し、マルテンサイトより軟質なベイニティックフェライトやパーライトの量を多くすること、つまり、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6を満たす必要がある。 Martensite area ratio / (Bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6
To ensure good material stability, reduce the amount of martensite that causes the material variation in the phase structure of the second phase, and increase the amount of bainitic ferrite and pearlite that are softer than martensite. That is, it is necessary to satisfy the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6.
良好な材質安定性を確保するために、第二相の相構成を、材質バラツキの要因となるマルテンサイトの量を低減し、マルテンサイトより軟質なベイニティックフェライトやパーライトの量を多くすること、つまり、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6を満たす必要がある。 Martensite area ratio / (Bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6
To ensure good material stability, reduce the amount of martensite that causes the material variation in the phase structure of the second phase, and increase the amount of bainitic ferrite and pearlite that are softer than martensite. That is, it is necessary to satisfy the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6.
なお、フェライト、ベイニティックフェライト、パーライト、マルテンサイト以外に、残留オーステナイトや焼戻しマルテンサイトやセメンタイト等の炭化物が生成する場合があるが、上記のフェライト、ベイニティックフェライト、パーライト、マルテンサイトの面積率が満足されていれば、本発明の目的を達成できる。
In addition to ferrite, bainitic ferrite, pearlite, and martensite, carbides such as retained austenite, tempered martensite, and cementite may be generated. The area of the above ferrite, bainitic ferrite, pearlite, and martensite If the rate is satisfied, the object of the present invention can be achieved.
また、本発明におけるフェライト、ベイニティックフェライト、パーライト、マルテンサイトの面積率とは、観察面積に占める各相の面積割合のことである。
Also, the area ratio of ferrite, bainitic ferrite, pearlite, and martensite in the present invention is the area ratio of each phase in the observation area.
本発明の高強度溶融亜鉛めっき鋼板は、上記成分組成と上記鋼組織を有する鋼板を下地鋼板とし、その上に溶融亜鉛めっきによるめっき皮膜、又は溶融亜鉛めっき後合金化処理を施しためっき皮膜を有する。
The high-strength hot-dip galvanized steel sheet of the present invention uses a steel sheet having the above component composition and the above steel structure as a base steel sheet, and a plated film obtained by hot-dip galvanizing, or a plated film that has been subjected to alloying treatment after hot-dip galvanizing. Have.
(3)次に製造条件について説明する。
(3) Next, manufacturing conditions will be described.
本発明の高強度溶融亜鉛めっき鋼板は、上記の成分組成範囲に適合した成分組成を有する鋼スラブを熱間圧延、酸洗し、またはさらに冷間圧延を行い、その後650℃以上の温度域まで5℃/s以上の平均加熱速度で加熱し、750~900℃の温度域で15~600s保持し、450~550℃の温度域に冷却し、該450~550℃の温度域にて10~200s保持し、次いで、溶融亜鉛めっきを施すことで製造する。
The high-strength hot-dip galvanized steel sheet according to the present invention is obtained by hot rolling, pickling, or further cold rolling a steel slab having a component composition suitable for the above-described component composition range, and thereafter up to a temperature range of 650 ° C. or higher. Heat at an average heating rate of 5 ° C./s or more, hold for 15 to 600 s in a temperature range of 750 to 900 ° C., cool to a temperature range of 450 to 550 ° C., and 10 to 10 in the temperature range of 450 to 550 ° C. It is manufactured by holding for 200 s and then applying hot dip galvanizing.
合金化処理を施す高強度溶融亜鉛めっき鋼板を製造するときは、溶融亜鉛めっき後、500~600℃の温度域において、下式を満たす条件で亜鉛めっきの合金化処理を施す。
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 When producing a high-strength hot-dip galvanized steel sheet subjected to alloying treatment, after hot-dip galvanizing, galvanizing alloying treatment is performed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 When producing a high-strength hot-dip galvanized steel sheet subjected to alloying treatment, after hot-dip galvanizing, galvanizing alloying treatment is performed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
以下、詳細に説明する。
The details will be described below.
上記の成分組成を有する鋼を、公知の方法により、溶製した後、分塊または連続鋳造を経てスラブとし、熱間圧延して熱延板にする。熱間圧延を行うに際しては、スラブを1100~1300℃に加熱し、最終仕上げ温度を850℃以上で熱間圧延を施し、400~650℃で鋼帯に巻き取ることが好ましい。巻き取り温度が650℃を超えた場合、熱延板中の炭化物が粗大化し、このような粗大化した炭化物は焼鈍時の均熱中に溶けきらないため、必要強度を得ることができない場合がある。その後、公知の方法で酸洗処理を行う。又は酸洗を行った後、さらに冷間圧延を行う。冷間圧延を行うに際しては、特にその条件を限定する必要はないが、30%以上の冷間圧下率で冷間圧延を施すことが好ましい。冷間圧下率が低いと、フェライトの再結晶が促進されず、未再結晶フェライトが残存し、延性と穴拡げ性が低下する場合があるためである。
The steel having the above composition is melted by a known method, and then slab is formed through a block or continuous casting, and is hot-rolled into a hot-rolled sheet. When performing hot rolling, it is preferable to heat the slab to 1100 to 1300 ° C., perform hot rolling at a final finishing temperature of 850 ° C. or higher, and wind it on a steel strip at 400 to 650 ° C. When the coiling temperature exceeds 650 ° C., the carbides in the hot-rolled sheet are coarsened, and such coarsened carbides cannot be melted during soaking at the time of annealing, so that the required strength may not be obtained. . Thereafter, pickling treatment is performed by a known method. Or after pickling, it cold-rolls further. When performing cold rolling, it is not necessary to specifically limit the conditions, but it is preferable to perform cold rolling at a cold reduction rate of 30% or more. If the cold rolling reduction is low, recrystallization of ferrite is not promoted, unrecrystallized ferrite remains, and ductility and hole expansibility may decrease.
酸洗した熱延板、または冷間圧延した鋼板に、以下の焼鈍を行った後、冷却し、その後溶融亜鉛めっきをする。
The hot-rolled sheet pickled or cold-rolled steel sheet is subjected to the following annealing, cooled, and then hot-dip galvanized.
650℃以上の温度域まで5℃/s以上の平均加熱速度で加熱
650℃以上の温度域までの平均加熱速度が5℃/s未満の場合、焼鈍中に微細で均一に分散したオーステナイト相が生成されず、最終組織のマルテンサイト面積率が増大し、良好な穴拡げ性の確保が困難である。また、通常よりも長い炉が必要となり、多大なエネルギー消費にともなうコスト増と生産効率の悪化を引き起こす。加熱炉としてDFF(Direct Fired Furnace)を用いることが好ましい。DFFによる急速加熱により、内部酸化層を形成させ、Si、Mn等の酸化物の鋼板最表層への濃化を防ぎ、良好なめっき性を確保するためである。 Heating at an average heating rate of 5 ° C./s or higher up to a temperature range of 650 ° C. or higher When the average heating rate up to a temperature range of 650 ° C. or higher is lower than 5 ° C./s, the austenite phase dispersed finely and uniformly during annealing It is not generated, the martensite area ratio of the final structure increases, and it is difficult to ensure good hole expansibility. In addition, a longer furnace than usual is required, which causes an increase in cost and deterioration in production efficiency due to a large amount of energy consumption. It is preferable to use DFF (Direct Fired Furnace) as the heating furnace. This is because an internal oxide layer is formed by rapid heating with DFF to prevent the oxide such as Si and Mn from being concentrated on the outermost surface layer of the steel sheet and to ensure good plating properties.
650℃以上の温度域までの平均加熱速度が5℃/s未満の場合、焼鈍中に微細で均一に分散したオーステナイト相が生成されず、最終組織のマルテンサイト面積率が増大し、良好な穴拡げ性の確保が困難である。また、通常よりも長い炉が必要となり、多大なエネルギー消費にともなうコスト増と生産効率の悪化を引き起こす。加熱炉としてDFF(Direct Fired Furnace)を用いることが好ましい。DFFによる急速加熱により、内部酸化層を形成させ、Si、Mn等の酸化物の鋼板最表層への濃化を防ぎ、良好なめっき性を確保するためである。 Heating at an average heating rate of 5 ° C./s or higher up to a temperature range of 650 ° C. or higher When the average heating rate up to a temperature range of 650 ° C. or higher is lower than 5 ° C./s, the austenite phase dispersed finely and uniformly during annealing It is not generated, the martensite area ratio of the final structure increases, and it is difficult to ensure good hole expansibility. In addition, a longer furnace than usual is required, which causes an increase in cost and deterioration in production efficiency due to a large amount of energy consumption. It is preferable to use DFF (Direct Fired Furnace) as the heating furnace. This is because an internal oxide layer is formed by rapid heating with DFF to prevent the oxide such as Si and Mn from being concentrated on the outermost surface layer of the steel sheet and to ensure good plating properties.
750~900℃の温度域で15~600s保持
750~900℃の温度域にて、具体的には、オーステナイト単相域、もしくはオーステナイトとフェライトの二相域で、15~600s保持する焼鈍を行う。焼鈍温度が750℃未満、保持時間が15s未満になると、鋼板中の硬質なセメンタイトが十分に溶解せず、穴拡げ性が低下し、さらに、所望のマルテンサイト面積率が得られないことから、延性が低下する。一方、焼鈍温度が900℃を超えると、オーステナイト粒の成長が著しく、冷却後の保持中に生じるベイナイト変態によるベイニティックフェライトの確保が困難となり、穴拡げ性が低下し、さらに、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)が0.6を超えるため、良好な材質安定性が得られない。また、保持時間が600sを超えると、オーステナイトが粗大化し、所望の強度確保が困難となり、また、多大なエネルギー消費にともなうコスト増を引き起こす場合がある。 Hold for 15 to 600 s in the temperature range of 750 to 900 ° C. In the temperature range of 750 to 900 ° C., specifically, anneal to hold for 15 to 600 s in the austenite single-phase region or the two-phase region of austenite and ferrite. . When the annealing temperature is less than 750 ° C. and the holding time is less than 15 s, the hard cementite in the steel sheet is not sufficiently dissolved, the hole expandability is lowered, and further, the desired martensite area ratio cannot be obtained. Ductility decreases. On the other hand, when the annealing temperature exceeds 900 ° C., the growth of austenite grains is remarkable, it becomes difficult to secure bainitic ferrite due to bainite transformation that occurs during holding after cooling, hole expansibility decreases, and the martensite area Since the ratio / (bainitic ferrite area ratio + pearlite area ratio) exceeds 0.6, good material stability cannot be obtained. On the other hand, if the holding time exceeds 600 s, austenite becomes coarse, and it becomes difficult to secure a desired strength, and it may cause an increase in cost due to a large energy consumption.
750~900℃の温度域にて、具体的には、オーステナイト単相域、もしくはオーステナイトとフェライトの二相域で、15~600s保持する焼鈍を行う。焼鈍温度が750℃未満、保持時間が15s未満になると、鋼板中の硬質なセメンタイトが十分に溶解せず、穴拡げ性が低下し、さらに、所望のマルテンサイト面積率が得られないことから、延性が低下する。一方、焼鈍温度が900℃を超えると、オーステナイト粒の成長が著しく、冷却後の保持中に生じるベイナイト変態によるベイニティックフェライトの確保が困難となり、穴拡げ性が低下し、さらに、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)が0.6を超えるため、良好な材質安定性が得られない。また、保持時間が600sを超えると、オーステナイトが粗大化し、所望の強度確保が困難となり、また、多大なエネルギー消費にともなうコスト増を引き起こす場合がある。 Hold for 15 to 600 s in the temperature range of 750 to 900 ° C. In the temperature range of 750 to 900 ° C., specifically, anneal to hold for 15 to 600 s in the austenite single-phase region or the two-phase region of austenite and ferrite. . When the annealing temperature is less than 750 ° C. and the holding time is less than 15 s, the hard cementite in the steel sheet is not sufficiently dissolved, the hole expandability is lowered, and further, the desired martensite area ratio cannot be obtained. Ductility decreases. On the other hand, when the annealing temperature exceeds 900 ° C., the growth of austenite grains is remarkable, it becomes difficult to secure bainitic ferrite due to bainite transformation that occurs during holding after cooling, hole expansibility decreases, and the martensite area Since the ratio / (bainitic ferrite area ratio + pearlite area ratio) exceeds 0.6, good material stability cannot be obtained. On the other hand, if the holding time exceeds 600 s, austenite becomes coarse, and it becomes difficult to secure a desired strength, and it may cause an increase in cost due to a large energy consumption.
450~550℃の温度域にて10~200s保持
前記の焼鈍を行った後、450~550℃の温度域に冷却し、該450~550℃の温度域に10~200s保持する。保持温度が550℃を超えると、または保持時間が10s未満になると、ベイナイト変態が促進せず、ベイニティックフェライトの面積率が1.0%未満になり、所望の穴拡げ性を得られない。また、保持温度が450℃未満、または保持時間が200sを超えると、第二相の大半がベイナイト変態の促進により生成した固溶炭素量の多いオーステナイトとベイニティックフェライトになり、所望の1.0%以上のパーライト面積率が得られず、かつ、硬質なマルテンサイト相の面積率が5.0%以上となり、良好な穴拡げ性と材質安定性が得られない。 Holding for 10 to 200 s in a temperature range of 450 to 550 ° C. After performing the above annealing, the temperature is cooled to a temperature range of 450 to 550 ° C. and held in the temperature range of 450 to 550 ° C. for 10 to 200 s. When the holding temperature exceeds 550 ° C. or the holding time is less than 10 s, the bainite transformation is not promoted, the area ratio of bainitic ferrite is less than 1.0%, and the desired hole expandability cannot be obtained. . When the holding temperature is less than 450 ° C. or the holding time exceeds 200 s, most of the second phase becomes austenite and bainitic ferrite having a large amount of dissolved carbon produced by promoting bainite transformation. A pearlite area ratio of 0% or more cannot be obtained, and an area ratio of the hard martensite phase is 5.0% or more, and good hole expansibility and material stability cannot be obtained.
前記の焼鈍を行った後、450~550℃の温度域に冷却し、該450~550℃の温度域に10~200s保持する。保持温度が550℃を超えると、または保持時間が10s未満になると、ベイナイト変態が促進せず、ベイニティックフェライトの面積率が1.0%未満になり、所望の穴拡げ性を得られない。また、保持温度が450℃未満、または保持時間が200sを超えると、第二相の大半がベイナイト変態の促進により生成した固溶炭素量の多いオーステナイトとベイニティックフェライトになり、所望の1.0%以上のパーライト面積率が得られず、かつ、硬質なマルテンサイト相の面積率が5.0%以上となり、良好な穴拡げ性と材質安定性が得られない。 Holding for 10 to 200 s in a temperature range of 450 to 550 ° C. After performing the above annealing, the temperature is cooled to a temperature range of 450 to 550 ° C. and held in the temperature range of 450 to 550 ° C. for 10 to 200 s. When the holding temperature exceeds 550 ° C. or the holding time is less than 10 s, the bainite transformation is not promoted, the area ratio of bainitic ferrite is less than 1.0%, and the desired hole expandability cannot be obtained. . When the holding temperature is less than 450 ° C. or the holding time exceeds 200 s, most of the second phase becomes austenite and bainitic ferrite having a large amount of dissolved carbon produced by promoting bainite transformation. A pearlite area ratio of 0% or more cannot be obtained, and an area ratio of the hard martensite phase is 5.0% or more, and good hole expansibility and material stability cannot be obtained.
その後、鋼板を通常の浴温のめっき浴中に浸入させて溶融亜鉛めっきを施し、ガスワイピングなどでめっき付着量を調整し、冷却することで、めっき層を合金化していない溶融亜鉛めっき鋼板を得る。
After that, the hot dip galvanized steel sheet in which the plated layer is not alloyed is cooled by infiltrating the steel sheet into a plating bath at a normal bath temperature and applying hot dip galvanizing. obtain.
合金化処理を施す溶融亜鉛めっき鋼板を製造するときは、溶融亜鉛めっきを施した後、さらに、500~600℃の温度域において、下式を満たす条件で亜鉛めっきの合金化処理を行う。
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 When producing a hot-dip galvanized steel sheet to be subjected to alloying treatment, after hot-dip galvanizing, galvanizing alloying treatment is further performed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 When producing a hot-dip galvanized steel sheet to be subjected to alloying treatment, after hot-dip galvanizing, galvanizing alloying treatment is further performed in a temperature range of 500 to 600 ° C. under conditions satisfying the following formula.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
exp[200/(400−T)]×ln(t)が0.45未満になると、合金化処理後の鋼組織にマルテンサイトが多く存在し、上記硬質なマルテンサイトが軟質なフェライトと隣接して、異相間に大きな硬度差が生じ、穴拡げ性が低下する。また、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)が0.6を超えるため、材質安定性が損なわれる。また、溶融亜鉛めっき層の付着性が悪くなる。
exp[200/(400−T)]×ln(t)が1.0超になると、未変態オーステナイトの殆どがセメンタイトもしくはパーライトに変態し、結果として所望の強度と延性のバランスが得られない。 When exp [200 / (400-T)] × ln (t) is less than 0.45, a lot of martensite is present in the steel structure after the alloying treatment, and the hard martensite is adjacent to soft ferrite. As a result, a large hardness difference occurs between the different phases, and the hole expansibility decreases. Further, since the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) exceeds 0.6, the material stability is impaired. Moreover, the adhesiveness of the hot dip galvanized layer is deteriorated.
When exp [200 / (400-T)] × ln (t) exceeds 1.0, most of the untransformed austenite is transformed into cementite or pearlite, and as a result, the desired balance between strength and ductility cannot be obtained.
exp[200/(400−T)]×ln(t)が1.0超になると、未変態オーステナイトの殆どがセメンタイトもしくはパーライトに変態し、結果として所望の強度と延性のバランスが得られない。 When exp [200 / (400-T)] × ln (t) is less than 0.45, a lot of martensite is present in the steel structure after the alloying treatment, and the hard martensite is adjacent to soft ferrite. As a result, a large hardness difference occurs between the different phases, and the hole expansibility decreases. Further, since the martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) exceeds 0.6, the material stability is impaired. Moreover, the adhesiveness of the hot dip galvanized layer is deteriorated.
When exp [200 / (400-T)] × ln (t) exceeds 1.0, most of the untransformed austenite is transformed into cementite or pearlite, and as a result, the desired balance between strength and ductility cannot be obtained.
500℃未満の温度域では、めっき層の合金化が促進されず、合金化溶融亜鉛めっき鋼板を得ることが難しい。また、600℃を超える温度域では、第二相の殆どがパーライトになり、所望のマルテンサイト面積率が得られず、強度と延性のバランスが低下する。
めっき層の合金化については、500~600℃の温度域において上記のexp[200/(400−T)]×ln(t)の条件を満たす本発明範囲であれば問題なく行うことができる。 In the temperature range below 500 ° C., alloying of the plating layer is not promoted, and it is difficult to obtain an galvannealed steel sheet. In the temperature range exceeding 600 ° C., most of the second phase becomes pearlite, the desired martensite area ratio cannot be obtained, and the balance between strength and ductility is lowered.
The alloying of the plating layer can be performed without any problem within the range of the present invention that satisfies the above condition of exp [200 / (400-T)] × ln (t) in a temperature range of 500 to 600 ° C.
めっき層の合金化については、500~600℃の温度域において上記のexp[200/(400−T)]×ln(t)の条件を満たす本発明範囲であれば問題なく行うことができる。 In the temperature range below 500 ° C., alloying of the plating layer is not promoted, and it is difficult to obtain an galvannealed steel sheet. In the temperature range exceeding 600 ° C., most of the second phase becomes pearlite, the desired martensite area ratio cannot be obtained, and the balance between strength and ductility is lowered.
The alloying of the plating layer can be performed without any problem within the range of the present invention that satisfies the above condition of exp [200 / (400-T)] × ln (t) in a temperature range of 500 to 600 ° C.
なお、本発明の製造方法における一連の熱処理においては、上述した温度範囲内であれば保持温度は一定である必要はなく、また冷却速度が冷却中に変化した場合においても規定した範囲内であれば本発明の趣旨を損なわない。また、熱履歴さえ満足されれば、鋼板はいかなる設備で熱処理を施されてもかまわない。加えて、熱処理後に形状矯正のため本発明の鋼板に調質圧延をすることも本発明の範囲に含まれる。なお、本発明では、鋼素材を通常の製鋼、鋳造、熱延の各工程を経て製造する場合を想定しているが、例えば薄手鋳造などにより熱延工程の一部もしくは全部を省略して製造する場合でもよい。
In the series of heat treatments in the production method of the present invention, the holding temperature does not need to be constant as long as it is within the above-mentioned temperature range, and even if the cooling rate changes during cooling, it may be within the specified range. Thus, the gist of the present invention is not impaired. Further, as long as the thermal history is satisfied, the steel sheet may be heat-treated by any equipment. In addition, temper rolling of the steel sheet of the present invention for shape correction after heat treatment is also included in the scope of the present invention. In the present invention, it is assumed that the steel material is manufactured through normal steelmaking, casting, and hot rolling processes, but the manufacturing process is performed by omitting part or all of the hot rolling process by thin casting, for example. You may do it.
図1、図2は、後述する実施例の本発明例である鋼AのNo.15、16、17(表2、表5)と比較例である鋼HのNo.18、19、20(表2、表5)について、TS、ELと焼鈍温度(T1)の関係を整理した図を示す。図1、図2より、本発明例の鋼Aは焼鈍温度の変化に伴なうTS、ELの変動が小さいのに対し、比較例の鋼HはTS、ELの変動が大きいことがわかる。
1 and 2 show No. of Steel A which is an example of the present invention in Examples described later. 15, 16, and 17 (Tables 2 and 5) and No. of Steel H as a comparative example. 18, 19, 20 (Table 2, Table 5) shows a diagram to organize TS, the relationship EL and annealing temperature (T 1). 1 and 2, it can be seen that the steel A of the present invention has a small variation in TS and EL accompanying changes in the annealing temperature, whereas the steel H of the comparative example has a large variation in TS and EL.
また、図3、図4は、後述する実施例の本発明例である鋼AのNo.21、22、23(表2、表5)と比較例である鋼HのNo.24、25、26(表2、表5)について、TS、ELと焼鈍後の冷却の平均保持温度(T2)の関係を整理した図を示す。図3、図4より、本発明例の鋼Aは平均保持温度の変化に伴なうTS、ELの変動が小さいのに対し、比較例の鋼HはTS、ELの変動が大きいことがわかる。
3 and 4 show the No. of steel A which is an example of the present invention in the examples described later. Nos. 21, 22, and 23 (Tables 2 and 5) and No. of Steel H as a comparative example. 24, 25, 26 (Table 2, Table 5) shows a diagram to organize TS, the relationship EL and the average holding temperature of the cooling after annealing (T 2). 3 and 4, it can be seen that the steel A of the present invention has a small variation in TS and EL accompanying changes in the average holding temperature, whereas the steel H of the comparative example has a large variation in TS and EL. .
表1に示す成分組成を有し、残部がFeおよび不可避的不純物よりなる鋼を転炉にて溶製し、連続鋳造法にてスラブとした。得られたスラブを1200℃に加熱後、870~920℃の仕上温度で板厚3.2mmまで熱間圧延を行い、520℃で巻き取った。次いで、得られた熱延板を酸洗し、一部は酸洗ままの熱延鋼板とし、一部はさらに冷間圧延を施し、冷延鋼板を製造した。次いで、上記により得られた熱延鋼板(酸洗後)および冷延鋼板を連続溶融亜鉛めっきラインにより、表2~表4に示す製造条件で、焼鈍処理を行い、溶融亜鉛めっき処理を施し、さらにめっき層の合金化処理を行い、溶融亜鉛めっき鋼板を得た。めっき付着量は片面あたり30~50g/m2とした。溶融亜鉛めっき処理を施した後に合金化処理を施さない溶融亜鉛めっき鋼板も一部作製した。
Steel having the component composition shown in Table 1 and the balance being Fe and inevitable impurities was melted in a converter and made into a slab by a continuous casting method. The obtained slab was heated to 1200 ° C., hot-rolled to a plate thickness of 3.2 mm at a finishing temperature of 870 to 920 ° C., and wound at 520 ° C. Subsequently, the obtained hot-rolled sheet was pickled, a part was made into a hot-rolled steel sheet as pickled, and a part was further cold-rolled to produce a cold-rolled steel sheet. Subsequently, the hot-rolled steel sheet (after pickling) and the cold-rolled steel sheet obtained as described above are subjected to an annealing treatment under the production conditions shown in Tables 2 to 4 by a continuous hot-dip galvanizing line, and hot-dip galvanizing treatment is performed. Furthermore, the alloying process of the plating layer was performed and the hot dip galvanized steel plate was obtained. The amount of plating was 30 to 50 g / m 2 per side. A part of a hot-dip galvanized steel sheet that was not subjected to alloying treatment after hot-dip galvanizing treatment was also produced.
得られた溶融亜鉛めっき鋼板に対して、フェライト、ベイニティックフェライト、パーライト、マルテンサイト相の面積率は、鋼板の圧延方向に平行な板厚断面を研磨後、3%ナイタールで腐食し、SEM(走査型電子顕微鏡)を用いて2000倍の倍率で10視野観察し、Media Cybernetics社のImage−Proを用いて求めた。その際、マルテンサイトと残留オーステナイトの区別が困難なため、得られた溶融亜鉛めっき鋼板に200℃で2時間の焼戻し処理を施し、その後、鋼板の圧延方向に平行な板厚断面の組織を上記の方法で観察し、上記の方法で求めた焼戻しマルテンサイト相の面積率をマルテンサイト相の面積率とした。また、残留オーステナイト相の体積率は、鋼板を板厚方向の1/4面まで研磨し、この板厚1/4面の回折X線強度により求めた。入射X線にはCoKα線を使用し、残留オーステナイト相の{111}、{200}、{220}、{311}面とフェライト相の{110}、{200}、{211}面のピークの積分強度の全ての組み合わせについて強度比を求め、これらの平均値を残留オーステナイト相の体積率とした。
The area ratio of ferrite, bainitic ferrite, pearlite, and martensite phase with respect to the obtained hot dip galvanized steel sheet was corroded with 3% nital after polishing the plate thickness section parallel to the rolling direction of the steel sheet, and SEM Ten fields of view were observed using a (scanning electron microscope) at a magnification of 2000 times, and obtained using Image-Pro of Media Cybernetics. At that time, since it is difficult to distinguish between martensite and retained austenite, the obtained hot-dip galvanized steel sheet was subjected to tempering treatment at 200 ° C. for 2 hours, and then the structure of the plate thickness section parallel to the rolling direction of the steel sheet was described above. The area ratio of the tempered martensite phase obtained by the above method was defined as the area ratio of the martensite phase. Further, the volume ratio of the retained austenite phase was obtained by polishing the steel plate to a ¼ surface in the plate thickness direction and diffracting X-ray intensity of the ¼ surface thickness. CoKα rays are used as incident X-rays, and the peaks of {111}, {200}, {220}, {311} planes of retained austenite phase and {110}, {200}, {211} planes of ferrite phase are used. Intensity ratios were obtained for all combinations of integrated intensities, and the average value of these ratios was taken as the volume ratio of the retained austenite phase.
また、引張試験は、引張方向が鋼板の圧延方向と直角方向となるようにサンプルを採取したJIS5号試験片を用いて、JIS Z 2241に準拠して行い、TS(引張強度)、EL(全伸び)を測定した。なお、本発明では、TS×EL≧19000MPa・%の場合を延性が良好と判定した。
In addition, the tensile test is performed in accordance with JIS Z 2241 using a JIS No. 5 test piece obtained by taking a sample so that the tensile direction is perpendicular to the rolling direction of the steel sheet, and TS (tensile strength), EL (all Elongation) was measured. In the present invention, it was determined that the ductility was good when TS × EL ≧ 19000 MPa ·%.
材質安定性は、(イ)焼鈍温度T1以外の条件が同じで焼鈍温度T1だけが異なる鋼板について、TS、ELの変動量を調査し、そのTS、ELの変動量から焼鈍温度変化20℃あたりの変動量(ΔTS、ΔEL)を求め、また(ロ)冷却後めっき浴浸漬までの平均保持温度T2以外の条件が同じで冷却後めっき浴浸漬までの平均保持温度T2だけが異なる鋼板について、TS、ELの変動量を調査し、そのTS、ELの変動量から冷却後めっき浴浸漬までの平均保持温度変化20℃あたりの変動量(ΔTS、ΔEL)を求め、各温度変化20℃当たりのTS変動量(ΔTS)、EL変動量(ΔEL)で評価した。
Material stability, (b) for only annealing temperatures T 1 condition the same annealing temperature T 1 of the non different steel, TS, to investigate the variation of EL, annealing temperature changes from the TS, the amount of variation of EL 20 variation amount per ° C. (.DELTA.TS, .DELTA.EL) the determined and (ii) only the average holding temperature T 2 to the average holding temperature T 2 than the condition are the same cooling after the plating bath immersion until the cooling after the plating bath immersion is different For the steel sheet, the amount of variation of TS and EL was investigated, and the amount of variation (ΔTS, ΔEL) per 20 ° C. of the average holding temperature change from the variation amount of TS and EL to the immersion in the plating bath after cooling was obtained, and each temperature change 20 Evaluation was made based on TS fluctuation amount (ΔTS) and EL fluctuation amount (ΔEL) per ° C.
また、以上により得られた溶融亜鉛めっき鋼板に対して、穴拡げ性(伸びフランジ性)を測定した。穴拡げ性(伸びフランジ性)は、日本鉄鋼連盟規格JFST1001に準拠して行った。得られた各鋼板を100mm×100mmに切断後、板厚2.0mm以上はクリアランス12%±1%で、板厚2.0mm未満はクリアランス12%±2%で、直径10mmの穴を打ち抜いた後、内径75mmのダイスを用いてしわ押さえ力9tonで抑えた状態で、60°円錐のポンチを穴に押し込んで亀裂発生限界における穴直径を測定し、下記の式から、限界穴広げ率λ(%)を求め、この限界穴広げ率の値から伸びフランジ性を評価した。
限界穴広げ率λ(%)={(Df−D0)/D0}×100
ただし、Dfは亀裂発生時の穴径(mm)、D0は初期穴径(mm)である。なお、本発明では、λ≧70(%)の場合を良好と判定した。 Moreover, the hole expansibility (stretch flangeability) was measured with respect to the hot dip galvanized steel sheet obtained by the above. The hole expandability (stretch flangeability) was performed in accordance with Japan Iron and Steel Federation Standard JFST1001. After each steel plate obtained was cut to 100 mm × 100 mm, a hole having a diameter of 10 mm or more was punched with a clearance of 12% ± 1% when the thickness was 2.0 mm or more, and a clearance of 12% ± 2% when the thickness was less than 2.0 mm. After that, using a die having an inner diameter of 75 mm and holding a wrinkle holding force of 9 ton, a 60 ° conical punch was pushed into the hole, and the hole diameter at the crack initiation limit was measured. %), And the stretch flangeability was evaluated from the value of the critical hole expansion rate.
Limit hole expansion ratio λ (%) = {(D f −D 0 ) / D 0 } × 100
However, D f hole diameter at crack initiation (mm), D 0 is the initial hole diameter (mm). In the present invention, the case of λ ≧ 70 (%) was determined to be good.
限界穴広げ率λ(%)={(Df−D0)/D0}×100
ただし、Dfは亀裂発生時の穴径(mm)、D0は初期穴径(mm)である。なお、本発明では、λ≧70(%)の場合を良好と判定した。 Moreover, the hole expansibility (stretch flangeability) was measured with respect to the hot dip galvanized steel sheet obtained by the above. The hole expandability (stretch flangeability) was performed in accordance with Japan Iron and Steel Federation Standard JFST1001. After each steel plate obtained was cut to 100 mm × 100 mm, a hole having a diameter of 10 mm or more was punched with a clearance of 12% ± 1% when the thickness was 2.0 mm or more, and a clearance of 12% ± 2% when the thickness was less than 2.0 mm. After that, using a die having an inner diameter of 75 mm and holding a wrinkle holding force of 9 ton, a 60 ° conical punch was pushed into the hole, and the hole diameter at the crack initiation limit was measured. %), And the stretch flangeability was evaluated from the value of the critical hole expansion rate.
Limit hole expansion ratio λ (%) = {(D f −D 0 ) / D 0 } × 100
However, D f hole diameter at crack initiation (mm), D 0 is the initial hole diameter (mm). In the present invention, the case of λ ≧ 70 (%) was determined to be good.
以上により得られた結果を表5~表7に示す。
Tables 5 to 7 show the results obtained as described above.
本発明例の高強度溶融亜鉛めっき鋼板は、いずれもTSが540MPa以上であり、λが70%以上で穴拡げ性に優れ、また、TS×EL≧19000MPa・%で強度と延性のバランスが高く、加工性に優れた高強度溶融亜鉛めっき鋼板であることがわかる。さらに、ΔTS、ΔELの値も小さく、材質安定性に優れた高強度溶融亜鉛めっき鋼板であることがわかる。一方、比較例では、延性、穴拡げ性のいずれか一つ以上が劣っているか、材質安定性が好ましくない。
Each of the high-strength hot-dip galvanized steel sheets of the present invention has a TS of 540 MPa or more, an excellent hole expansibility when λ is 70% or more, and a high balance of strength and ductility when TS × EL ≧ 19000 MPa ·%. It can be seen that this is a high-strength hot-dip galvanized steel sheet with excellent workability. Furthermore, it can be seen that the values of ΔTS and ΔEL are small, and the steel sheet is a high-strength hot-dip galvanized steel sheet excellent in material stability. On the other hand, in the comparative example, one or more of ductility and hole expansibility are inferior, or material stability is not preferable.
本発明の高強度溶融亜鉛めっき鋼板は、540MPa以上の引張強度TSを有し、かつ、高延性と高穴拡げ性を有し、さらに材質安定性にも優れる。本発明の高強度溶融亜鉛めっき鋼板を、例えば、自動車構造部材に適用することにより車体軽量化による燃費改善を図ることができ、産業上の利用価値は非常に大きい。
The high-strength hot-dip galvanized steel sheet of the present invention has a tensile strength TS of 540 MPa or more, has high ductility and high hole expansibility, and is excellent in material stability. By applying the high-strength hot-dip galvanized steel sheet of the present invention to, for example, automobile structural members, fuel efficiency can be improved by reducing the weight of the vehicle body, and the industrial utility value is very large.
Claims (8)
- 成分組成は、質量%でC:0.04%以上0.13%以下、Si:0.7%以上2.3%以下、Mn:0.8%以上2.0%以下、P:0.1%以下、S:0.01%以下、Al:0.1%以下、N:0.008%以下を含有し、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、75%以上のフェライト相と、1.0%以上のベイニティックフェライト相と、1.0%以上10.0%以下のパーライト相を有し、さらに、マルテンサイト相の面積率が1.0%以上5.0%未満で、かつ、マルテンサイト面積率/(ベイニティックフェライト面積率+パーライト面積率)≦0.6を満たすことを特徴とする材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 The component composition is C: 0.04% to 0.13%, Si: 0.7% to 2.3%, Mn: 0.8% to 2.0%, P: 0.00% by mass. 1% or less, S: 0.01% or less, Al: 0.1% or less, N: 0.008% or less, with the balance being Fe and inevitable impurities, and the steel structure is 75 in area ratio. % Of ferrite phase, 1.0% or more of bainitic ferrite phase, 1.0% or more and 10.0% or less of pearlite phase, and the area ratio of martensite phase is 1.0% More than 5.0% and satisfying martensite area ratio / (bainitic ferrite area ratio + pearlite area ratio) ≦ 0.6 Galvanized steel sheet.
- さらに、成分組成として、質量%で、Cr:0.05%以上1.0%以下、V:0.005%以上0.5%以下、Mo:0.005%以上0.5%以下、Ni:0.05%以上1.0%以下、Cu:0.05%以上1.0%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ことを特徴とする請求項1に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 Furthermore, as a component composition, in mass%, Cr: 0.05% or more and 1.0% or less, V: 0.005% or more and 0.5% or less, Mo: 0.005% or more and 0.5% or less, Ni 2. The material stability according to claim 1, comprising at least one element selected from: 0.05% to 1.0%, Cu: 0.05% to 1.0%. High-strength hot-dip galvanized steel sheet with excellent workability and workability.
- さらに、成分組成として、質量%で、Ti:0.01%以上0.1%以下、Nb:0.01%以上0.1%以下、B:0.0003%以上0.0050%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ことを特徴とする請求項1または2に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 Furthermore, as a component composition, in mass%, Ti: 0.01% or more and 0.1% or less, Nb: 0.01% or more and 0.1% or less, B: 0.0003% or more and 0.0050% or less The high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to claim 1 or 2, comprising at least one element selected from the group consisting of:
- さらに、成分組成として、質量%で、Ca:0.001%以上0.005%以下、REM:0.001%以上0.005%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ことを特徴とする請求項1~3のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 Furthermore, as a component composition, it contains at least one element selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005% in mass%. The high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of claims 1 to 3.
- さらに、成分組成として、質量%で、Ta:0.001%以上0.010%以下、Sn:0.002%以上0.2%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ことを特徴とする請求項1~4のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 Furthermore, as a component composition, it contains at least one element selected from Ta: 0.001% to 0.010% and Sn: 0.002% to 0.2% by mass%. The high-strength hot-dip galvanized steel sheet excellent in material stability and workability according to any one of claims 1 to 4.
- さらに、成分組成として、質量%で、Sb:0.002%以上0.2%以下を含有することを特徴とする請求項1~5のいずれかに記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板。 Furthermore, as a component composition, Sb: 0.002% or more and 0.2% or less is contained in mass%, and excellent in material stability and workability according to any one of claims 1 to 5 High strength hot dip galvanized steel sheet.
- 請求項1~6のいずれかに記載の成分組成を有する鋼スラブを、熱間圧延、酸洗し、またはさらに冷間圧延し、その後650℃以上の温度域まで5℃/s以上の平均加熱速度で加熱し、750~900℃の温度域で15~600s保持し、450~550℃の温度域に冷却した後、該450~550℃の温度域で10~200s保持し、次いで、溶融亜鉛めっきを施すことを特徴とする材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。 A steel slab having the component composition according to any one of claims 1 to 6 is hot-rolled, pickled, or further cold-rolled, and then heated at an average temperature of 5 ° C / s or higher to a temperature range of 650 ° C or higher. Heat at a speed, hold for 15 to 600 s in a temperature range of 750 to 900 ° C, cool to a temperature range of 450 to 550 ° C, hold for 10 to 200 s in the temperature range of 450 to 550 ° C, and then molten zinc A method for producing a high-strength hot-dip galvanized steel sheet excellent in material stability and workability, characterized by performing plating.
- 溶融亜鉛めっきを施した後、500~600℃の温度域において下式を満たす条件で亜鉛めっきの合金化処理を施すことを特徴とする請求項7に記載の材質安定性と加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
0.45≦exp[200/(400−T)]×ln(t)≦1.0
但し、
T:500~600℃の温度域での平均保持温度(℃)
t:500~600℃の温度域の保持時間(s)
exp(X)、ln(X)は、それぞれXの指数関数、自然対数を示す。 8. The material stability and workability of claim 7, wherein after the hot dip galvanizing, the alloying treatment of the galvanizing is performed under a condition satisfying the following formula in a temperature range of 500 to 600 ° C. A method for producing high-strength hot-dip galvanized steel sheets.
0.45 ≦ exp [200 / (400−T)] × ln (t) ≦ 1.0
However,
T: Average holding temperature in the temperature range of 500 to 600 ° C (° C)
t: Holding time in the temperature range of 500 to 600 ° C. (s)
exp (X) and ln (X) denote the exponential function and natural logarithm of X, respectively.
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- 2011-01-18 WO PCT/JP2011/051151 patent/WO2011090180A1/en active Application Filing
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EP2402470A1 (en) * | 2009-02-25 | 2012-01-04 | JFE Steel Corporation | High-strength hot-dip galvanized steel plate of excellent workability and manufacturing method therefor |
EP2402470A4 (en) * | 2009-02-25 | 2017-04-26 | JFE Steel Corporation | High-strength hot-dip galvanized steel plate of excellent workability and manufacturing method therefor |
EP2623629A4 (en) * | 2010-09-30 | 2016-08-31 | Jfe Steel Corp | High-strength hot-dip galvanized steel sheet having excellent fatigue characteristics and manufacturing method therefor |
JP2012177175A (en) * | 2011-02-28 | 2012-09-13 | Jfe Steel Corp | Low-yield-ratio and high-strength cold rolled steel sheet having excellent elongation and excellent stretch flangeability and method for producing same |
WO2012165661A1 (en) * | 2011-06-01 | 2012-12-06 | Jfeスチール株式会社 | Process for producing high-strength hot-dip galvanized steel sheet with excellent material-quality stability, processability, and deposit appearance |
US9340859B2 (en) | 2011-06-01 | 2016-05-17 | Jfe Steel Corporation | Method for manufacturing high strength galvanized steel sheet having excellent stability of mechanical properties, formability, and coating appearance |
EP2762583A4 (en) * | 2011-09-30 | 2015-12-02 | Nippon Steel & Sumitomo Metal Corp | High-strength hot-dip galvanized steel sheet having excellent delayed fracture resistance, and method for producing same |
WO2013179497A1 (en) * | 2012-06-01 | 2013-12-05 | Jfeスチール株式会社 | Low yield ratio high-strength cold-rolled steel sheet with excellent elongation and stretch flange formability, and manufacturing method thereof |
KR20150004430A (en) * | 2012-06-01 | 2015-01-12 | 제이에프이 스틸 가부시키가이샤 | Low yield ratio high-strength cold-rolled steel sheet with excellent elongation and stretch flange formability, and manufacturing method thereof |
KR101674283B1 (en) | 2012-06-01 | 2016-11-08 | 제이에프이 스틸 가부시키가이샤 | High strength cold-rolled steel sheet with low yield ratio having excellent elongation and stretch flangeability, and method for manufacturing the same |
EP2886674A4 (en) * | 2012-08-15 | 2016-11-30 | Nippon Steel & Sumitomo Metal Corp | Steel sheet for hot pressing use, method for producing same, and hot press steel sheet member |
US10570470B2 (en) | 2012-08-15 | 2020-02-25 | Nippon Steel Corporation | Steel sheet for hot stamping, method of manufacturing the same, and hot stamped steel sheet member |
Also Published As
Publication number | Publication date |
---|---|
TWI433961B (en) | 2014-04-11 |
JP2011168877A (en) | 2011-09-01 |
EP2527482B1 (en) | 2019-12-25 |
JP5786317B2 (en) | 2015-09-30 |
EP2527482A4 (en) | 2017-04-05 |
EP2527482A1 (en) | 2012-11-28 |
TW201139731A (en) | 2011-11-16 |
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