EP2604715B1 - Method for manufacturing a high-strength cold-rolled steel sheet having excellent formability and crashworthiness - Google Patents
Method for manufacturing a high-strength cold-rolled steel sheet having excellent formability and crashworthiness Download PDFInfo
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- EP2604715B1 EP2604715B1 EP10855912.1A EP10855912A EP2604715B1 EP 2604715 B1 EP2604715 B1 EP 2604715B1 EP 10855912 A EP10855912 A EP 10855912A EP 2604715 B1 EP2604715 B1 EP 2604715B1
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- Prior art keywords
- temperature
- steel sheet
- martensite
- rolling
- rolled steel
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- 239000010960 cold rolled steel Substances 0.000 title claims description 20
- 238000000034 method Methods 0.000 title claims description 19
- 238000004519 manufacturing process Methods 0.000 title claims description 18
- 229910000831 Steel Inorganic materials 0.000 claims description 54
- 239000010959 steel Substances 0.000 claims description 54
- 229910001566 austenite Inorganic materials 0.000 claims description 53
- 229910000734 martensite Inorganic materials 0.000 claims description 50
- 238000001816 cooling Methods 0.000 claims description 40
- 230000000717 retained effect Effects 0.000 claims description 30
- 238000010438 heat treatment Methods 0.000 claims description 23
- 230000009466 transformation Effects 0.000 claims description 17
- 229910000859 α-Fe Inorganic materials 0.000 claims description 17
- 238000000137 annealing Methods 0.000 claims description 15
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical compound [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 claims description 14
- 238000005098 hot rolling Methods 0.000 claims description 12
- 239000000203 mixture Substances 0.000 claims description 9
- 229910052799 carbon Inorganic materials 0.000 claims description 8
- 238000005097 cold rolling Methods 0.000 claims description 8
- 229910052802 copper Inorganic materials 0.000 claims description 7
- 229910052750 molybdenum Inorganic materials 0.000 claims description 7
- 229910052759 nickel Inorganic materials 0.000 claims description 7
- 239000012535 impurity Substances 0.000 claims description 6
- 229910052742 iron Inorganic materials 0.000 claims description 6
- 229910052758 niobium Inorganic materials 0.000 claims description 6
- 229910052710 silicon Inorganic materials 0.000 claims description 6
- 229910052720 vanadium Inorganic materials 0.000 claims description 6
- 229910052782 aluminium Inorganic materials 0.000 claims description 5
- 229910052748 manganese Inorganic materials 0.000 claims description 3
- 229910052698 phosphorus Inorganic materials 0.000 claims description 3
- 229910052717 sulfur Inorganic materials 0.000 claims description 3
- 230000000052 comparative effect Effects 0.000 description 32
- 230000000694 effects Effects 0.000 description 18
- 238000005096 rolling process Methods 0.000 description 17
- 230000007423 decrease Effects 0.000 description 11
- 239000011572 manganese Substances 0.000 description 10
- 238000003303 reheating Methods 0.000 description 9
- PXHVJJICTQNCMI-UHFFFAOYSA-N nickel Substances [Ni] PXHVJJICTQNCMI-UHFFFAOYSA-N 0.000 description 8
- 239000011651 chromium Substances 0.000 description 7
- 239000010949 copper Substances 0.000 description 7
- 238000005728 strengthening Methods 0.000 description 7
- 230000003247 decreasing effect Effects 0.000 description 6
- 239000010955 niobium Substances 0.000 description 6
- 229910001562 pearlite Inorganic materials 0.000 description 6
- 239000010936 titanium Substances 0.000 description 6
- 239000011575 calcium Substances 0.000 description 5
- 229910052761 rare earth metal Inorganic materials 0.000 description 5
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 4
- 238000010521 absorption reaction Methods 0.000 description 4
- 238000005266 casting Methods 0.000 description 4
- 229910001567 cementite Inorganic materials 0.000 description 4
- 229910052804 chromium Inorganic materials 0.000 description 4
- 239000013078 crystal Substances 0.000 description 4
- 239000000463 material Substances 0.000 description 4
- 230000002411 adverse Effects 0.000 description 3
- 229910001563 bainite Inorganic materials 0.000 description 3
- 238000009749 continuous casting Methods 0.000 description 3
- 230000000593 degrading effect Effects 0.000 description 3
- 230000006866 deterioration Effects 0.000 description 3
- 238000005516 engineering process Methods 0.000 description 3
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 3
- 238000005461 lubrication Methods 0.000 description 3
- 238000001556 precipitation Methods 0.000 description 3
- 238000001953 recrystallisation Methods 0.000 description 3
- 239000002436 steel type Substances 0.000 description 3
- 238000009864 tensile test Methods 0.000 description 3
- 229910052719 titanium Inorganic materials 0.000 description 3
- LFQSCWFLJHTTHZ-UHFFFAOYSA-N Ethanol Chemical compound CCO LFQSCWFLJHTTHZ-UHFFFAOYSA-N 0.000 description 2
- 229910001335 Galvanized steel Inorganic materials 0.000 description 2
- XUIMIQQOPSSXEZ-UHFFFAOYSA-N Silicon Chemical compound [Si] XUIMIQQOPSSXEZ-UHFFFAOYSA-N 0.000 description 2
- 229910000794 TRIP steel Inorganic materials 0.000 description 2
- XAGFODPZIPBFFR-UHFFFAOYSA-N aluminium Chemical compound [Al] XAGFODPZIPBFFR-UHFFFAOYSA-N 0.000 description 2
- 230000003749 cleanliness Effects 0.000 description 2
- 239000011248 coating agent Substances 0.000 description 2
- 238000000576 coating method Methods 0.000 description 2
- 238000005336 cracking Methods 0.000 description 2
- 239000000446 fuel Substances 0.000 description 2
- 239000008397 galvanized steel Substances 0.000 description 2
- 238000005246 galvanizing Methods 0.000 description 2
- 229910052751 metal Inorganic materials 0.000 description 2
- 239000002184 metal Substances 0.000 description 2
- 150000001247 metal acetylides Chemical class 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 238000010791 quenching Methods 0.000 description 2
- 230000000171 quenching effect Effects 0.000 description 2
- 239000010703 silicon Substances 0.000 description 2
- 238000002791 soaking Methods 0.000 description 2
- 239000000243 solution Substances 0.000 description 2
- 239000013585 weight reducing agent Substances 0.000 description 2
- UXHQLGLGLZKHTC-CUNXSJBXSA-N 4-[(3s,3ar)-3-cyclopentyl-7-(4-hydroxypiperidine-1-carbonyl)-3,3a,4,5-tetrahydropyrazolo[3,4-f]quinolin-2-yl]-2-chlorobenzonitrile Chemical compound C1CC(O)CCN1C(=O)C1=CC=C(C=2[C@@H]([C@H](C3CCCC3)N(N=2)C=2C=C(Cl)C(C#N)=CC=2)CC2)C2=N1 UXHQLGLGLZKHTC-CUNXSJBXSA-N 0.000 description 1
- ZOXJGFHDIHLPTG-UHFFFAOYSA-N Boron Chemical compound [B] ZOXJGFHDIHLPTG-UHFFFAOYSA-N 0.000 description 1
- OYPRJOBELJOOCE-UHFFFAOYSA-N Calcium Chemical compound [Ca] OYPRJOBELJOOCE-UHFFFAOYSA-N 0.000 description 1
- VYZAMTAEIAYCRO-UHFFFAOYSA-N Chromium Chemical compound [Cr] VYZAMTAEIAYCRO-UHFFFAOYSA-N 0.000 description 1
- RYGMFSIKBFXOCR-UHFFFAOYSA-N Copper Chemical compound [Cu] RYGMFSIKBFXOCR-UHFFFAOYSA-N 0.000 description 1
- 229910000885 Dual-phase steel Inorganic materials 0.000 description 1
- PWHULOQIROXLJO-UHFFFAOYSA-N Manganese Chemical compound [Mn] PWHULOQIROXLJO-UHFFFAOYSA-N 0.000 description 1
- ZOKXTWBITQBERF-UHFFFAOYSA-N Molybdenum Chemical compound [Mo] ZOKXTWBITQBERF-UHFFFAOYSA-N 0.000 description 1
- GRYLNZFGIOXLOG-UHFFFAOYSA-N Nitric acid Chemical compound O[N+]([O-])=O GRYLNZFGIOXLOG-UHFFFAOYSA-N 0.000 description 1
- 239000004186 Penicillin G benzathine Substances 0.000 description 1
- OAICVXFJPJFONN-UHFFFAOYSA-N Phosphorus Chemical compound [P] OAICVXFJPJFONN-UHFFFAOYSA-N 0.000 description 1
- NINIDFKCEFEMDL-UHFFFAOYSA-N Sulfur Chemical compound [S] NINIDFKCEFEMDL-UHFFFAOYSA-N 0.000 description 1
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 1
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 description 1
- 238000002441 X-ray diffraction Methods 0.000 description 1
- 229910045601 alloy Inorganic materials 0.000 description 1
- 239000000956 alloy Substances 0.000 description 1
- 230000003466 anti-cipated effect Effects 0.000 description 1
- 229910052796 boron Inorganic materials 0.000 description 1
- 229910052791 calcium Inorganic materials 0.000 description 1
- -1 carbon nitrides Chemical class 0.000 description 1
- 239000003795 chemical substances by application Substances 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 238000005304 joining Methods 0.000 description 1
- 230000000670 limiting effect Effects 0.000 description 1
- 239000003595 mist Substances 0.000 description 1
- 239000011733 molybdenum Substances 0.000 description 1
- GUCVJGMIXFAOAE-UHFFFAOYSA-N niobium atom Chemical compound [Nb] GUCVJGMIXFAOAE-UHFFFAOYSA-N 0.000 description 1
- 229910017604 nitric acid Inorganic materials 0.000 description 1
- 239000003921 oil Substances 0.000 description 1
- 230000003647 oxidation Effects 0.000 description 1
- 238000007254 oxidation reaction Methods 0.000 description 1
- 239000011574 phosphorus Substances 0.000 description 1
- 238000005498 polishing Methods 0.000 description 1
- 238000004080 punching Methods 0.000 description 1
- 230000002829 reductive effect Effects 0.000 description 1
- 238000007670 refining Methods 0.000 description 1
- 239000011347 resin Substances 0.000 description 1
- 229920005989 resin Polymers 0.000 description 1
- 229920006395 saturated elastomer Polymers 0.000 description 1
- 238000005204 segregation Methods 0.000 description 1
- 230000035939 shock Effects 0.000 description 1
- 239000006104 solid solution Substances 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
- 239000011593 sulfur Substances 0.000 description 1
- 239000002344 surface layer Substances 0.000 description 1
- 230000003746 surface roughness Effects 0.000 description 1
- 239000000725 suspension Substances 0.000 description 1
- 238000005496 tempering Methods 0.000 description 1
- LEONUFNNVUYDNQ-UHFFFAOYSA-N vanadium atom Chemical compound [V] LEONUFNNVUYDNQ-UHFFFAOYSA-N 0.000 description 1
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 1
- 230000004584 weight gain Effects 0.000 description 1
- 235000019786 weight gain Nutrition 0.000 description 1
Classifications
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—Hardening; Quenching with or without subsequent tempering
- C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
-
- 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
- C21D2211/001—Austenite
-
- 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
-
- 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
-
- 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
- C21D9/02—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for springs
Definitions
- the present invention relates to a method for manufacturing a high-strength cold rolled steel sheet having excellent formability for use in structural parts and suspension parts mainly used in the automobile industry field.
- DP steel ferrite-martensite dual phase steel
- TRIP steel that utilizes the transformation-induced plasticity of retained austenite
- Patent Literature 1 discloses a method for manufacturing a high-strength steel sheet having good formability, with which high ductility is achieved by adding large quantities of Si and thereby reliably obtaining retained austenite.
- the stretch flangeability is an indicator of formability during flange-forming through expanding holes already made, and is a property as important as the elongation property required for high-strength steel sheets.
- Patent Literature 2 discloses a method for manufacturing a cold rolled steel sheet having good stretch flangeability with which the stretch flangeability is improved by forming a ferrite-tempered martensite multi-phase microstructure by conducting quenching and tempering after annealing and soaking.
- this technology although high stretch flangeability is achieved, the elongation is low.
- PTL 3 describes a method for manufacturing a high tensile strength galvanized steel sheet with excellent formability, comprising the steps of subjecting a slab having an elemental composition comprising, in terms of % by mass, 0.05 to 0.3 % of C, 0.01 to 2.5 % of Si, 0.5 to 3.5 % of Mn, 0.003 to 0.100 % of P, 0.02 % or less of S, 0.010 to 1.5 % of Al and 0.007 % or less of N, the remainder being Fe and unavoidable impurities, to hot rolling and cold rolling thereby making a cold rolled steel sheet, subjecting the cold rolled steel sheet to annealing including steps of heating and maintaining the steel sheet in a temperature range from 750 to 950°C for 10 seconds or more, cooling the steel sheet from 750°C to a temperature range from (Ms point - 100°C) to (Ms point - 200°C) at an average cooling rate of 10°C/s or more, and reheating and maintaining the steel sheet in
- PTL 4 describes a method for manufacturing a high-strength galvanized steel sheet with excellent formability, comprising the steps of hot rolling a slab that contains, in terms of % by mass, 0.05 to 0.3 % C, 0.01 to 2.5 % Si, 0.5 to 3.5 % of Mn, 0.003 to 0.100 % or less of P, 0.02 % or less of S and 0.010 to 1.5 % of Al, the total of Si and Al being 0.5 to 2.5 %, the remainder being iron and incidental impurities, to form a steel sheet; in continuous annealing, heating the steel sheet to a temperature in the range of 750 to 900°C at an average heating rate of at least 10°C/s in the temperature range of 500°C to an A 1 transformation point, holding that temperature for at least 10 seconds, cooling the steel sheet from 750°C to a temperature in the range of (Ms point - 100°C) to (Ms point - 200°C) at an average cooling rate of at least 10°C/s
- the present invention has been made by addressing the problems described above and an object of the present invention is to provide a method for manufacturing a high-strength cold rolled steel sheet having excellent ductility and stretch flangeability.
- the inventors of the invention of the present application have conducted extensive studies from the points of view of steel sheet composition and microstructure so as to address the problems described above and manufacture a high-strength cold rolled steel sheet having excellent ductility and stretch flangeability. As a result, they have found that when the steel has alloy elements adequately controlled, intensively cooled to a 150 to 350°C temperature range during cooling from the soaking temperature in the annealing process, and reheated, a microstructure containing 20% or more ferrite and 10 to 60% tempered martensite in terms of area ratio and 3 to 15% retained austenite in terms of volume ratio can be obtained and high ductility and stretch flangeability can be achieved.
- the ductility improves due to a TRIP effect of the retained austenite.
- the martensite generated by transformation of retained austenite under application of strain becomes very hard and as a result exhibits a hardness significantly different from that of the main phase ferrite, thereby degrading the stretch flangeability.
- this steel sheet microstructure can exhibit high formability and improved crashworthiness.
- the present invention provides a method for manufacturing a high-strength cold rolled steel sheet having excellent formability and crashworthiness, the method comprising hot-rolling and cold-rolling a slab having a composition comprising, on a mass% basis, C: 0.05 to 0.3%, Si: 0.3 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100%, S: 0.02% or less, Al: 0.010 to 0.5%, optionally at least one element selected from Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.005 to 2.00%, and Cu: 0.005 to 2.00%, optionally one or both of Ti: 0.01 to 0.20% and Nb: 0.01 to 0.20%, optionally B: 0.0002 to 0.005%, optionally one or both of Ca: 0.001 to 0.005% and REM: 0.001 to 0.005%, and balance being iron and unavoidable impurities, to manufacture a cold rolled steel sheet and
- the present invention a high-strength cold rolled steel sheet having excellent formability is obtained.
- the present invention achieves advantageous effects such as realizing both weight reduction and improved crash safety of automobiles and greatly contributing to improving performance of automobile bodies.
- Carbon (C) is an element that stabilizes austenite and promotes generation of phases other than ferrite. Thus, carbon is needed to increase the steel sheet strength, generate a multiphase structure, and improve the TS-EL balance. At a C content less than 0.05%, it is difficult to reliably obtain phases other than ferrite even when the production conditions are optimized and TS ⁇ EL decreases as a result. At a C content exceeding 0.3%, hardening of welded portions and heat-affected zones is significant, and mechanical properties of the welded portions are deteriorated. Thus, the C content is within the range of 0.05 to 0.3% and preferably 0.08 to 0.15%.
- Silicon (Si) is an element effective for strengthening the steel. Silicon is also a ferrite-generating element, suppresses C from becoming concentrated and forming carbides in the austenite, and thus serves to accelerate generation of retained austenite.
- the Si content is less than 0.3%, the effects of addition are low. Thus, the lower limit is 0.3%. Excessive addition deteriorates the surface quality and weldability. Thus, the Si content is 2.5% or less.
- the Si content is preferably in the range of 0.7 to 2.0%.
- Manganese (Mn) is an element effective for strengthening the steel and accelerates generation of low-temperature transformation-forming phase such as tempered martensite. Such an effect is observed at a Mn content of 0.5% or more. However, when the Mn content exceeds 3.5%, the second phase fraction increases excessively, the ductility deterioration of ferrite due to solid solution strengthening becomes significant, and formability is degraded. Accordingly, the Mn content is within the range of 0.5 to 3.5% and preferably in the range of 1.5 to 3.0%.
- Phosphorus (P) is an element effective for strengthening the steel and this effect is achieved at a P content of 0.003% or more.
- P content is in the range of 0.003% to 0.100%.
- S Sulfur
- the S content is preferably as low as possible but is limited to 0.02% or less from the production cost point of view.
- Aluminum (Al) acts as a deoxidizing agent and is an element effective for cleanliness of the steel. Aluminum is preferably added in the deoxidizing process. When the Al content is less than 0.01%, the effect of addition is little and thus the lower limit is 0.01%. However, addition of large quantities of Al increases the risk of slab cracking during continuous casting and decreases the productivity. Thus, the upper limit of the Al content is 0.5%.
- the high-strength cold rolled steel sheet manufactured by the method of the present invention contains the above-described components as the basic components and the balance is iron and unavoidable impurities. However, the following components can be adequately contained according to the desired properties.
- Chromium (Cr), molybdenum (Mo), vanadium (V), nickel (Ni), and copper (Cu) suppress generation of pearlite during cooling from the annealing temperature, promote generation of low-temperature transformation-forming phases, and effectively serves to strengthen the steel.
- Cr, Mo, V, Ni, and Cu are contained.
- the content of each of Cr, Mo, V, Ni, and Cu exceeds 2.00%, the effect is saturated and the cost will rise.
- the Cr, Mo, V, Ni, and Cu contents are each in the range of 0.005 to 2.00%.
- Titanium (Ti) and niobium (Nb) form carbon nitrides and have an effect of strengthening the steel by precipitation. Such effects are observed at a content of 0.01% or more for each element.
- Ti and Nb are contained in amounts exceeding 0.20%, excessive strengthening occurs and the ductility is decreased.
- the Ti and Nb contents are each within the range of 0.01 to 0.20%.
- B Boron
- Ca Ca
- REM rare earth element
- the area fraction of the ferrite is less than 20%, TS ⁇ EL decreases.
- the area fraction of ferrite is limited to 20% or more and preferably 50% or more.
- Tempered martensite is a ferrite-cementite multiphase having a high dislocation density and is obtained by heating martensite to a temperature equal to or lower than Ac 1 transformation point and preferably to a temperature lower than Ac 1 transformation point. Tempered martensite effectively strengthens the steel.
- the microstructure obtained by heating martensite to a temperature exceeding Ac 1 transformation point is a microstructure that does not contain cementite in ferrite and is fundamentally different from the tempered martensite intended in the present invention.
- the tempered martensite Compared to martensite, the tempered martensite has less adverse effects on stretch flangeability and is a phase effective for reliably obtaining the strength without significantly decreasing the stretch flangeability.
- the area fraction of the tempered martensite is less than 10%, it becomes difficult to reliably obtain the strength.
- the area fraction exceeds 60%, TS ⁇ EL is decreased.
- the area fraction of the martensite is limited to 10 to 60%.
- Martensite effectively increases the strength of the steel but significantly decreases the stretch flangeability once the area fraction of the martensite exceeds 10%. Thus, the area fraction of the martensite is limited to 0 to 10%.
- volume fraction of retained austenite 3 to 15%
- Retained austenite not only contributes to strengthening of the steel but also effectively improves TS ⁇ EL of the steel. Such effects are achieved at a volume fraction of 3% or more. When the volume fraction of the retained austenite exceeds 15%, the stretch flangeability is decreased. Accordingly, the volume fraction of the retained austenite is limited to 3 to 15%.
- Average crystal grain diameter of low-temperature transformation-forming phases constituted by martensite, tempered martensite, and retained austenite 3 ⁇ m or less
- Low-temperature transformation-forming phases constituted by martensite, tempered martensite, and retained austenite effectively improve the crashworthiness.
- finely dispersing the low-temperature transformation-forming phases improves the crashworthiness, and this effect becomes notable when the average crystal grain diameter of the low-temperature transformation-forming phases is 3 ⁇ m or less. Accordingly, the average crystal grain diameter of the low-temperature transformation-forming phases is limited to 3 ⁇ m or less.
- the phases other than ferrite, tempered martensite, martensite, and retained austenite may include pearlite and bainite but such phases do not present problem as long as the above-described phase structure is satisfied.
- the pearlite is preferably 3% or less from the view points of ductility and stretch flangeability.
- a steel having a composition controlled as described above is melted in a converter or the like and formed into a slab by continuous casting or the like.
- This steel is hot-rolled, cold-rolled, and continuously annealed.
- the manufacturing methods regarding casting, hot-rolling, and cold-rolling are not particularly limited but preferable manufacturing methods are described below.
- the steel slab used is preferably manufactured by continuous casting in order to prevent macrosegregation of the components but an ingot casting technique or a thin slab casting technique may be employed.
- an energy-saving process such as hot direct rolling or direct rolling which involves sending the hot slab to a heating furnace without cooling the slab to room temperature or which involves rolling the slab immediately after a short period of heat retention may be employed without any difficulty.
- the slab heating temperature is preferably low from the viewpoint of energy. At a heating temperature less than 1100°C, carbides cannot be sufficiently dissolved or the risks of troubles during hot-rolling increases due to an increased rolling load. In order to prevent the increase in scale loss attributable to oxidation weight gain, the slab heating temperature is preferably 1300°C or less.
- a sheet bar heater that heats the sheet bar may be employed.
- Finishing temperature Ar 3 transformation point or more.
- the finishing temperature is less than the Ar 3 transformation point, ferrite and austenite are generated during rolling, and a band-like microstructure readily occurs in the steel sheet. Such a band-like microstructure remains after cold rolling and annealing, may generate anisotropy in the material properties, and may decrease the formability. Accordingly, the finishing temperature is preferably equal to or higher than Ar 3 transformation point.
- the coiling temperature is preferably in the range of 450 to 700°C.
- part or all of the finish rolling may be conducted by lubrication rolling.
- Lubrication rolling is effective from the viewpoints of uniform steel sheet shape and material homogeneity.
- the coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10.
- Preferable is a continuous rolling process of joining sheet bars next to each other and continuously finish-rolling the sheet bars.
- the continuous rolling process is also preferable from the viewpoint of operation stability of hot rolling.
- the oxidized scales on the surface of the hot-rolled steel sheet are preferably removed by pickling and the steel sheet is cold-rolled to form a cold-rolled steel sheet having a particular thickness.
- the pickling conditions and the cold rolling conditions are not particularly limited and typical conditions may be used.
- the reduction of cold rolling is preferably 40% or more.
- Average heating rate from 500°C to Ac 1 transformation point 10°C/s or more
- the average heating rate in the recrystallization temperature zone, 500°C to Ac 1 transformation point, of the steel of the present invention is 10°C/s or more, recrystallization during heating is suppressed, austenite generated at Ac 1 transformation temperature or higher becomes finer, and the microstructure after annealing and cooling becomes finer.
- the average grain diameter of the low-temperature transformation-forming phase can be reduced to 3 ⁇ m or less.
- the average heating rate from 500°C to Ac 1 transformation point is limited to 10°C/s or more and more preferably 20°C/s or more.
- the heating temperature is less than 750°C or the holding time is less than 10 seconds, generation of austenite during annealing is insufficient and a sufficient amount of low-temperature transformation-forming phases cannot be reliably obtained after annealing and cooling.
- the upper limits of the holding temperature and the holding time are not particularly defined, the effects saturate and the cost will increase when the holding temperature is 900°C or more and the holding time is 600 seconds or more. Accordingly, the holding temperature is preferably less than 900°C and the holding time is preferably less than 600 seconds.
- the cooling rate from 750°C is less than 10°C/s, pearlite is generated and TS ⁇ EL and stretch flangeability are degraded.
- the cooling rate from 750°C is limited to 10°C/s or more.
- the temperature condition of ending the cooling is one of the most crucial conditions of this technology. At the time cooling is stopped, part of austenite transforms into martensite and the rest forms untransformed austenite. When reheated, plated and alloyed, and cooled to room temperature, martensite turns into tempered martensite and untransformed austenite transforms into retained austenite or martensite.
- controlling the temperature of ending the cooling determines the final area fractions of the martensite, the retained austenite, and the tempered martensite.
- the temperature of ending the cooling is higher than 350°C, martensite transformation at the time cooling is stopped is insufficient and the amount of untransformed austenite is large, thereby ultimately generating excessive amounts of martensite or retained austenite and degrading the stretch flangeability.
- the temperature of ending the cooling is lower than 150°C, most of austenite transforms into martensite during cooling, the amount of untransformed austenite decreases, and 3% or more of retained austenite is not obtained. Accordingly, the temperature of ending the cooling is set within the range of 150 to 350°C.
- any cooling method such as gas jet cooling, mist cooling, water cooling, or metal quenching, may be employed as long as the target cooling rate and cooling end temperature are achieved.
- the martensite generated during cooling is tempered and forms tempered martensite.
- the stretch flangeability is improved, the untransformed austenite that did not transform into martensite during cooling is stabilized, and 3% or more of retained austenite is obtained at the final stage, thereby improving the ductility.
- the reheating temperature is less than 350°C, the martensite is not sufficiently tempered and the austenite is not sufficiently stabilized, thereby degrading stretch flangeability and ductility. If the reheating temperature exceeds 600°C, untransformed austenite at the time cooling is stopped transforms into pearlite and 3% or more of retained austenite cannot be obtained at the final stage. Accordingly, the heating temperature is limited to 350 to 600°C.
- the reheating temperature is set within the range of 350 to 600°C and the holding time within that temperature range is limited to 10 to 600 seconds.
- the annealed steel sheet may be subjected to temper rolling to correct shape, adjust surface roughness, etc. Moreover, treatment such as resin or oil/fat coating and various other coating may be performed.
- a steel having the composition shown in Table 1 and balance being Fe and unavoidable impurities was melted in a converter and continuously casted into a slab.
- the slab is hot-rolled to a thickness of 3.0 mm.
- the hot rolling conditions were as follows: finishing temperature: 900°C, cooling rate after rolling: 10°C/s, and coiling temperature: 600°C. Then the hot-rolled steel sheet was pickled and cold-rolled to a thickness of 1.2 mm to manufacture a cold rolled steel sheet.
- the cold rolled steel sheet was annealed under the conditions described in Table 2 by using a continuous annealing line.
- the cross-sectional microstructure of the steel sheet was observed by exposing the microstructure by using a 3% nital solution (3% nitric acid + ethanol), observing the position 1/4 of the thickness in the depth direction by using a scanning electron microscope, and conducting an image processing of a picture of the microstructure taken to determine the fraction of the ferrite phase (the image processing can be performed by using commercially available image processing software).
- the area fractions of the martensite and tempered martensite were determined by taking SEM photographs of adequate magnification, e.g., about 1000 to 3000 magnification, depending on the fineness of the microstructure and then determining the quantity by using image processing software.
- the average grain diameter of the low-temperature transformation-forming phase was determined by dividing the area of the low-temperature transformation-forming phases in the observed area by the number of the low-temperature transformation-forming phases, determining the average area therefrom, and raising the average to the power of 1/2.
- the volume ratio of the retained austenite was determined by polishing the steel sheet to a surface 1/4 in the thickness direction and measuring X-ray diffraction intensity of the 1/4 thickness surface.
- a MoK ⁇ line was used as the incident X ray, the intensity ratios were determined for all combinations of the integrated intensities of peaks of ⁇ 111 ⁇ , ⁇ 200 ⁇ , ⁇ 220 ⁇ , and ⁇ 311 ⁇ faces of the retained austenite phase and the ⁇ 110 ⁇ , ⁇ 200 ⁇ , and ⁇ 211 ⁇ faces of the ferrite phase, and the average value was assumed to be the volume fraction of the retained austenite.
- the tensile property was determined by using a JIS No. 5 specimen sampled from the steel sheet in such a manner that the tensile direction was orthogonal to the rolling direction, conducting a tensile test according to JIS Z2241 to measure TS (tensile strength) and EL (elongation), and determining the strength-elongation balance value represented by the product of the strength and elongation (TS ⁇ EL).
- the hole expanding ratio ⁇ was measured as an indicator for evaluating the stretch flangeability.
- the hole expanding ratio ⁇ was determined by conducting a hole expanding test according to the Japan Iron and Steel Federation standard JFST1001 and determining the ratio from the initial diameter (10 mm ⁇ ) of the hole upon punching and the diameter of hole at the time the crack at the hole edge penetrated the sheet upon hole expanding.
- the shock absorption property was determined by using a specimen 5 mm in width and 7 mm in length sampled from the steel sheet in a direction orthogonal to the rolling direction, conducting a tensile test at a strain rate of 2000/s, and integrating the stress-true strain curve obtained by the tensile test within the range of 0 to 10% to calculate the absorption energy (refer to Tetsu-to-Hagane, 83 (1997) p. 748 ).
- the steel sheets of the examples of the present invention have excellent strength, ductility, and stretch flangeability, i.e., TS ⁇ EL of 22000 MPa ⁇ % or more and ⁇ of 70% or more.
- the steel sheets of comparative examples outside the range of the present invention did not achieve excellent strength, ductility, and stretch flangeability unlike the steel sheets of the examples of the present invention since TS ⁇ EL was less than 22000 MPa ⁇ % and/or ⁇ was less than 70%.
- the ratio of the absorption energy to TS is 0.063 or more, thereby achieving excellent crashworthiness.
- the present invention can contribute to weight reduction and decreasing the fuel consumption of automobiles by providing a high-strength cold rolled steel sheet having excellent formability and crashworthiness.
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Description
- The present invention relates to a method for manufacturing a high-strength cold rolled steel sheet having excellent formability for use in structural parts and suspension parts mainly used in the automobile industry field.
- In recent years, from the viewpoint of preserving global environment, improving fuel efficiency of automobiles has become a key issue. There has been a trend toward increasing the strength of automobile body materials to achieve thickness reduction and decrease the weight of car bodies. However, increasing the strength of steel sheets leads to a decrease in ductility, i.e., a decrease in formability and workability. Thus, development of materials that have both high strength and high formability has been anticipated.
- To fulfill such a need, various multi-phase cold rolled steel sheets including ferrite-martensite dual phase steel (hereinafter referred to as "DP steel") and TRIP steel that utilizes the transformation-induced plasticity of retained austenite have been developed.
- For example, Patent Literature 1 discloses a method for manufacturing a high-strength steel sheet having good formability, with which high ductility is achieved by adding large quantities of Si and thereby reliably obtaining retained austenite.
- However, although DP steel and TRIP steel have good elongation properties, they have poor stretch flangeability. The stretch flangeability is an indicator of formability during flange-forming through expanding holes already made, and is a property as important as the elongation property required for high-strength steel sheets.
- Patent Literature 2 discloses a method for manufacturing a cold rolled steel sheet having good stretch flangeability with which the stretch flangeability is improved by forming a ferrite-tempered martensite multi-phase microstructure by conducting quenching and tempering after annealing and soaking. However, according to this technology, although high stretch flangeability is achieved, the elongation is low.
- According to existing technologies, cold-rolled steel sheets having good elongation property and stretch flangeability have not been obtained.
- PTL 3 describes a method for manufacturing a high tensile strength galvanized steel sheet with excellent formability, comprising the steps of subjecting a slab having an elemental composition comprising, in terms of % by mass, 0.05 to 0.3 % of C, 0.01 to 2.5 % of Si, 0.5 to 3.5 % of Mn, 0.003 to 0.100 % of P, 0.02 % or less of S, 0.010 to 1.5 % of Al and 0.007 % or less of N, the remainder being Fe and unavoidable impurities, to hot rolling and cold rolling thereby making a cold rolled steel sheet, subjecting the cold rolled steel sheet to annealing including steps of heating and maintaining the steel sheet in a temperature range from 750 to 950°C for 10 seconds or more, cooling the steel sheet from 750°C to a temperature range from (Ms point - 100°C) to (Ms point - 200°C) at an average cooling rate of 10°C/s or more, and reheating and maintaining the steel sheet in a temperature range from 350 to 600°C for 1 to 600 seconds, and then subjecting the annealed steel sheet to galvanizing treatment.
- PTL 4 describes a method for manufacturing a high-strength galvanized steel sheet with excellent formability, comprising the steps of hot rolling a slab that contains, in terms of % by mass, 0.05 to 0.3 % C, 0.01 to 2.5 % Si, 0.5 to 3.5 % of Mn, 0.003 to 0.100 % or less of P, 0.02 % or less of S and 0.010 to 1.5 % of Al, the total of Si and Al being 0.5 to 2.5 %, the remainder being iron and incidental impurities, to form a steel sheet; in continuous annealing, heating the steel sheet to a temperature in the range of 750 to 900°C at an average heating rate of at least 10°C/s in the temperature range of 500°C to an A1 transformation point, holding that temperature for at least 10 seconds, cooling the steel sheet from 750°C to a temperature in the range of (Ms point - 100°C) to (Ms point - 200°C) at an average cooling rate of at least 10°C/s, reheating the steel sheet to a temperature in the range of 350 to 600°C, and holding that temperature for 10 to 600 seconds; and galvanizing the steel sheet.
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- PTL 1: Japanese Unexamined Patent Application Publication No.
2-101117 - PTL 2: Japanese Unexamined Patent Application Publication No.
2004-256872 - PTL 3:
WO 2009/054539 A1 - PTL 4:
WO 2009/096344 A1 - The present invention has been made by addressing the problems described above and an object of the present invention is to provide a method for manufacturing a high-strength cold rolled steel sheet having excellent ductility and stretch flangeability.
- The inventors of the invention of the present application have conducted extensive studies from the points of view of steel sheet composition and microstructure so as to address the problems described above and manufacture a high-strength cold rolled steel sheet having excellent ductility and stretch flangeability. As a result, they have found that when the steel has alloy elements adequately controlled, intensively cooled to a 150 to 350°C temperature range during cooling from the soaking temperature in the annealing process, and reheated, a microstructure containing 20% or more ferrite and 10 to 60% tempered martensite in terms of area ratio and 3 to 15% retained austenite in terms of volume ratio can be obtained and high ductility and stretch flangeability can be achieved.
- Typically, when retained austenite is present, the ductility improves due to a TRIP effect of the retained austenite. However, it is known that the martensite generated by transformation of retained austenite under application of strain becomes very hard and as a result exhibits a hardness significantly different from that of the main phase ferrite, thereby degrading the stretch flangeability.
- However, according to the composition and microstructure of the present invention, both high ductility and high stretch flangeability are achieved simultaneously. Although the exact reason why high stretch flangeability is achieved despite the presence of retained austenite is unclear, it is presumed that co-existence of the retained austenite and tempered martensite reduces the adverse effects of the retained austenite on the stretch flangeability.
- It has also been found that when the average crystal grain diameter of the low-temperature transformation-forming phase constituted by martensite, tempered martensite, and retained austenite is 3 µm or less, this steel sheet microstructure can exhibit high formability and improved crashworthiness.
- The present invention has been made based on the findings described above and is summarized as follows.
- The present invention provides a method for manufacturing a high-strength cold rolled steel sheet having excellent formability and crashworthiness, the method comprising hot-rolling and cold-rolling a slab having a composition comprising, on a mass% basis, C: 0.05 to 0.3%, Si: 0.3 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100%, S: 0.02% or less, Al: 0.010 to 0.5%, optionally at least one element selected from Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.005 to 2.00%, and Cu: 0.005 to 2.00%, optionally one or both of Ti: 0.01 to 0.20% and Nb: 0.01 to 0.20%, optionally B: 0.0002 to 0.005%, optionally one or both of Ca: 0.001 to 0.005% and REM: 0.001 to 0.005%, and balance being iron and unavoidable impurities, to manufacture a cold rolled steel sheet and continuously annealing the cold rolled sheet, wherein, during the continuous annealing, the steel sheet is heated with an average heating rate in the range of 500°C to Ac1 transformation point being 10°C/s or more, held at a temperature of 750°C or more for 10 seconds or more, cooled from 750°C to a temperature range of 150 to 350°C at a cooling rate of 10°C/s or more on average, heated to a temperature of 350 to 600°C, held thereat for 10 to 600 seconds, and cooled to room temperature.
- According to the present invention a high-strength cold rolled steel sheet having excellent formability is obtained. The present invention achieves advantageous effects such as realizing both weight reduction and improved crash safety of automobiles and greatly contributing to improving performance of automobile bodies.
- The present invention will now be described in detail.
- The reasons for limiting the steel composition to those described above are first described. Note that the meaning of % regarding components is mass% unless otherwise noted.
- Carbon (C) is an element that stabilizes austenite and promotes generation of phases other than ferrite. Thus, carbon is needed to increase the steel sheet strength, generate a multiphase structure, and improve the TS-EL balance. At a C content less than 0.05%, it is difficult to reliably obtain phases other than ferrite even when the production conditions are optimized and TS × EL decreases as a result. At a C content exceeding 0.3%, hardening of welded portions and heat-affected zones is significant, and mechanical properties of the welded portions are deteriorated. Thus, the C content is within the range of 0.05 to 0.3% and preferably 0.08 to 0.15%.
- Silicon (Si) is an element effective for strengthening the steel. Silicon is also a ferrite-generating element, suppresses C from becoming concentrated and forming carbides in the austenite, and thus serves to accelerate generation of retained austenite. When the Si content is less than 0.3%, the effects of addition are low. Thus, the lower limit is 0.3%. Excessive addition deteriorates the surface quality and weldability. Thus, the Si content is 2.5% or less. The Si content is preferably in the range of 0.7 to 2.0%.
- Manganese (Mn) is an element effective for strengthening the steel and accelerates generation of low-temperature transformation-forming phase such as tempered martensite. Such an effect is observed at a Mn content of 0.5% or more. However, when the Mn content exceeds 3.5%, the second phase fraction increases excessively, the ductility deterioration of ferrite due to solid solution strengthening becomes significant, and formability is degraded. Accordingly, the Mn content is within the range of 0.5 to 3.5% and preferably in the range of 1.5 to 3.0%.
- Phosphorus (P) is an element effective for strengthening the steel and this effect is achieved at a P content of 0.003% or more. When P is contained exceeding 0.100%, brittleness is induced by grain segregation and crashworthiness is deteriorated. Accordingly, the P content is in the range of 0.003% to 0.100%.
- Sulfur (S) forms inclusions such as MnS and causes deterioration of crashworthiness and cracking along the metal flow of the welded portion. Thus, the S content is preferably as low as possible but is limited to 0.02% or less from the production cost point of view.
- Aluminum (Al) acts as a deoxidizing agent and is an element effective for cleanliness of the steel. Aluminum is preferably added in the deoxidizing process. When the Al content is less than 0.01%, the effect of addition is little and thus the lower limit is 0.01%. However, addition of large quantities of Al increases the risk of slab cracking during continuous casting and decreases the productivity. Thus, the upper limit of the Al content is 0.5%.
- The high-strength cold rolled steel sheet manufactured by the method of the present invention contains the above-described components as the basic components and the balance is iron and unavoidable impurities. However, the following components can be adequately contained according to the desired properties.
- Chromium (Cr), molybdenum (Mo), vanadium (V), nickel (Ni), and copper (Cu) suppress generation of pearlite during cooling from the annealing temperature, promote generation of low-temperature transformation-forming phases, and effectively serves to strengthen the steel. Such effects are obtained when 0.005% or more of at least one of Cr, Mo, V, Ni, and Cu is contained. However, when the content of each of Cr, Mo, V, Ni, and Cu exceeds 2.00%, the effect is saturated and the cost will rise. Thus, the Cr, Mo, V, Ni, and Cu contents are each in the range of 0.005 to 2.00%.
- Titanium (Ti) and niobium (Nb) form carbon nitrides and have an effect of strengthening the steel by precipitation. Such effects are observed at a content of 0.01% or more for each element. In contrast, when Ti and Nb are contained in amounts exceeding 0.20%, excessive strengthening occurs and the ductility is decreased. Thus, the Ti and Nb contents are each within the range of 0.01 to 0.20%.
- Boron (B) suppresses generation of ferrite from austenite grain boundaries and increases the strength. Such effects are obtained at a B content of 0.0002% or more. However, the effects saturate when the B content exceeds 0.005% and the cost will rise. Accordingly, the B content is within the range of 0.0002 to 0.005%.
- Calcium (Ca) and a rare earth element (REM) improve the formability through sulfide morphology control. If needed, one or both of Ca and REM may be contained at an amount of 0.001% or more each. However, since excessive addition may adversely affect the cleanliness, the amount of each element is limited to 0.005% or less.
- The microstructure of the steel will now be described.
- When the area fraction of the ferrite is less than 20%, TS × EL decreases. Thus, the area fraction of ferrite is limited to 20% or more and preferably 50% or more.
- Tempered martensite is a ferrite-cementite multiphase having a high dislocation density and is obtained by heating martensite to a temperature equal to or lower than Ac1 transformation point and preferably to a temperature lower than Ac1 transformation point. Tempered martensite effectively strengthens the steel. The microstructure obtained by heating martensite to a temperature exceeding Ac1 transformation point is a microstructure that does not contain cementite in ferrite and is fundamentally different from the tempered martensite intended in the present invention.
- Compared to martensite, the tempered martensite has less adverse effects on stretch flangeability and is a phase effective for reliably obtaining the strength without significantly decreasing the stretch flangeability. When the area fraction of the tempered martensite is less than 10%, it becomes difficult to reliably obtain the strength. When the area fraction exceeds 60%, TS × EL is decreased. Thus, the area fraction of the martensite is limited to 10 to 60%.
- Martensite effectively increases the strength of the steel but significantly decreases the stretch flangeability once the area fraction of the martensite exceeds 10%. Thus, the area fraction of the martensite is limited to 0 to 10%.
- Retained austenite not only contributes to strengthening of the steel but also effectively improves TS × EL of the steel. Such effects are achieved at a volume fraction of 3% or more. When the volume fraction of the retained austenite exceeds 15%, the stretch flangeability is decreased. Accordingly, the volume fraction of the retained austenite is limited to 3 to 15%.
- Low-temperature transformation-forming phases constituted by martensite, tempered martensite, and retained austenite effectively improve the crashworthiness. In particular, finely dispersing the low-temperature transformation-forming phases improves the crashworthiness, and this effect becomes notable when the average crystal grain diameter of the low-temperature transformation-forming phases is 3 µm or less. Accordingly, the average crystal grain diameter of the low-temperature transformation-forming phases is limited to 3 µm or less.
- The phases other than ferrite, tempered martensite, martensite, and retained austenite may include pearlite and bainite but such phases do not present problem as long as the above-described phase structure is satisfied. However, the pearlite is preferably 3% or less from the view points of ductility and stretch flangeability.
- A steel having a composition controlled as described above is melted in a converter or the like and formed into a slab by continuous casting or the like. This steel is hot-rolled, cold-rolled, and continuously annealed. The manufacturing methods regarding casting, hot-rolling, and cold-rolling are not particularly limited but preferable manufacturing methods are described below.
- The steel slab used is preferably manufactured by continuous casting in order to prevent macrosegregation of the components but an ingot casting technique or a thin slab casting technique may be employed. In addition to an existing method of cooling the manufactured steel slab to room temperature and then reheating the slab, an energy-saving process such as hot direct rolling or direct rolling which involves sending the hot slab to a heating furnace without cooling the slab to room temperature or which involves rolling the slab immediately after a short period of heat retention may be employed without any difficulty.
- The slab heating temperature is preferably low from the viewpoint of energy. At a heating temperature less than 1100°C, carbides cannot be sufficiently dissolved or the risks of troubles during hot-rolling increases due to an increased rolling load. In order to prevent the increase in scale loss attributable to oxidation weight gain, the slab heating temperature is preferably 1300°C or less.
- In order to avoid troubles during hot-rolling despite the decreased slab heating temperature, a sheet bar heater that heats the sheet bar may be employed.
- When the finishing temperature is less than the Ar3 transformation point, ferrite and austenite are generated during rolling, and a band-like microstructure readily occurs in the steel sheet. Such a band-like microstructure remains after cold rolling and annealing, may generate anisotropy in the material properties, and may decrease the formability. Accordingly, the finishing temperature is preferably equal to or higher than Ar3 transformation point.
- When the coiling temperature is less than 450°C, the control of the coiling temperature is difficult and temperature nonuniformity may occur, thereby causing problems such as deterioration of cold-rolling properties. When the coiling temperature exceeds 700°C, problems such as decarburization in the base iron surface layer may occur. Thus, the coiling temperature is preferably in the range of 450 to 700°C.
- In the hot-rolling process of the present invention, in order to decrease the rolling load during hot rolling, part or all of the finish rolling may be conducted by lubrication rolling. Lubrication rolling is effective from the viewpoints of uniform steel sheet shape and material homogeneity. Note that the coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10. Preferable is a continuous rolling process of joining sheet bars next to each other and continuously finish-rolling the sheet bars. The continuous rolling process is also preferable from the viewpoint of operation stability of hot rolling.
- Next, the oxidized scales on the surface of the hot-rolled steel sheet are preferably removed by pickling and the steel sheet is cold-rolled to form a cold-rolled steel sheet having a particular thickness. The pickling conditions and the cold rolling conditions are not particularly limited and typical conditions may be used. The reduction of cold rolling is preferably 40% or more.
- When the average heating rate in the recrystallization temperature zone, 500°C to Ac1 transformation point, of the steel of the present invention is 10°C/s or more, recrystallization during heating is suppressed, austenite generated at Ac1 transformation temperature or higher becomes finer, and the microstructure after annealing and cooling becomes finer. As a result, the average grain diameter of the low-temperature transformation-forming phase can be reduced to 3 µm or less.
- When the average heating rate is less than 10°C/s, α recrystallization occurs during heating and strain introduced into ferrite is released and thus the sufficient refining of grains cannot be achieved. Thus, the average heating rate from 500°C to Ac1 transformation point is limited to 10°C/s or more and more preferably 20°C/s or more.
- When the heating temperature is less than 750°C or the holding time is less than 10 seconds, generation of austenite during annealing is insufficient and a sufficient amount of low-temperature transformation-forming phases cannot be reliably obtained after annealing and cooling. Although the upper limits of the holding temperature and the holding time are not particularly defined, the effects saturate and the cost will increase when the holding temperature is 900°C or more and the holding time is 600 seconds or more. Accordingly, the holding temperature is preferably less than 900°C and the holding time is preferably less than 600 seconds.
- When the cooling rate from 750°C is less than 10°C/s, pearlite is generated and TS × EL and stretch flangeability are degraded. Thus, the cooling rate from 750°C is limited to 10°C/s or more. The temperature condition of ending the cooling is one of the most crucial conditions of this technology. At the time cooling is stopped, part of austenite transforms into martensite and the rest forms untransformed austenite. When reheated, plated and alloyed, and cooled to room temperature, martensite turns into tempered martensite and untransformed austenite transforms into retained austenite or martensite. When the temperature of ending the cooling from annealing is low, the amount of martensite generated during cooling increases and the amount of the untransformed austenite decreases. Thus, controlling the temperature of ending the cooling determines the final area fractions of the martensite, the retained austenite, and the tempered martensite.
- When the temperature of ending the cooling is higher than 350°C, martensite transformation at the time cooling is stopped is insufficient and the amount of untransformed austenite is large, thereby ultimately generating excessive amounts of martensite or retained austenite and degrading the stretch flangeability. When the temperature of ending the cooling is lower than 150°C, most of austenite transforms into martensite during cooling, the amount of untransformed austenite decreases, and 3% or more of retained austenite is not obtained. Accordingly, the temperature of ending the cooling is set within the range of 150 to 350°C. As for the cooling method, any cooling method such as gas jet cooling, mist cooling, water cooling, or metal quenching, may be employed as long as the target cooling rate and cooling end temperature are achieved.
- When the steel is held in the temperature range of 350 to 600°C for 10 seconds or more after being cooled to a temperature range of 150 to 350°C, the martensite generated during cooling is tempered and forms tempered martensite. As a result, the stretch flangeability is improved, the untransformed austenite that did not transform into martensite during cooling is stabilized, and 3% or more of retained austenite is obtained at the final stage, thereby improving the ductility.
- Although the detailed mechanism of stabilization of the untransformed austenite by reheating and holding is not clear, it is presumed that carbon diffuses from martensite, in which dissolved C is oversaturated, into untransformed austenite, thereby increasing the C concentration in the untransformed austenite and stabilizing the austenite. During this process, if the precipitation of cementite in the martensite occurs faster than diffusion of carbon, the concentration of C in the untransformed austenite becomes insufficient. Thus, it is important to delay the cementite precipitation and this requires addition of 0.3% or more of Si.
- If the reheating temperature is less than 350°C, the martensite is not sufficiently tempered and the austenite is not sufficiently stabilized, thereby degrading stretch flangeability and ductility. If the reheating temperature exceeds 600°C, untransformed austenite at the time cooling is stopped transforms into pearlite and 3% or more of retained austenite cannot be obtained at the final stage. Accordingly, the heating temperature is limited to 350 to 600°C.
- If the holding time is less than 10 seconds, the austenite is not sufficiently stabilized. If the holding time exceeds 600 seconds, untransformed austenite at the time the cooling is stopped transforms into bainite and 3% or more of retained austenite cannot be obtained at the final stage. Accordingly, the reheating temperature is set within the range of 350 to 600°C and the holding time within that temperature range is limited to 10 to 600 seconds.
- The annealed steel sheet may be subjected to temper rolling to correct shape, adjust surface roughness, etc. Moreover, treatment such as resin or oil/fat coating and various other coating may be performed.
- A steel having the composition shown in Table 1 and balance being Fe and unavoidable impurities was melted in a converter and continuously casted into a slab. The slab is hot-rolled to a thickness of 3.0 mm. The hot rolling conditions were as follows: finishing temperature: 900°C, cooling rate after rolling: 10°C/s, and coiling temperature: 600°C. Then the hot-rolled steel sheet was pickled and cold-rolled to a thickness of 1.2 mm to manufacture a cold rolled steel sheet.
- The cold rolled steel sheet was annealed under the conditions described in Table 2 by using a continuous annealing line.
- The cross-sectional microstructure, tensile properties, and stretch flangeability of the resulting steel sheet were investigated. The results are shown in Table 3.
[Table 1] (mass%) Steel type C Si Mn P S Al N Cr Mo V Ni Cu Ti Nb B Ca REM A 0.10 1.2 2.3 0.020 0.003 0.033 0.003 Example B 0.07 1.7 2.0 0.025 0.003 0.036 0.004 0.30 Example C 0.18 1.0 1.6 0.013 0.005 0.028 0.005 0.4 Example D 0.25 1.5 1.4 0.008 0.006 0.031 0.003 0.05 Example E 0.08 0.5 2.2 0.007 0.003 0.030 0.002 0.2 0.4 Exam ple F 0.12 1.1 1.9 0.007 0.002 0.400 0.001 0.05 Example G 0.14 1.5 2.3 0.014 0.001 0.042 0.003 0.04 Example H 0.10 0.9 1.9 0.021 0.005 0.015 0.004 0.02 0.001 Example I 0.08 1.2 2.5 0.006 0.004 0.026 0.002 0.004 Example J 0.09 2.0 1.8 0.012 0.003 0.028 0.005 0.002 Example K 0.04 1.3 1.8 0.013 0.002 0.022 0.002 Comparative Example L 0.17 0.6 4.0 0.022 0.001 0.036 0.002 Comparative Example M 0.10 1.1 0.3 0.007 0.003 0.029 0.002 Comparative Example Note: Underlined items are outside the range of the present invention. [Table 2] No. Steel type AC1 transformation point Average heating rate from 500°C to Ac1 Maximum temperature Holding time Average cooling rate Temperature after cooling Reheating temperature Holding time after reheating °C °C/s °C Sec °C/s °C °C Sec 1 A 721 15 830 60 50 200 400 80 Example 2 A 15 810 60 50 100 420 80 Comparative Example 3 B 740 20 850 90 80 180 430 60 Example 4 B 20 720 60 80 250 430 60 Comparative Example 5 C 734 5 820 90 30 160 450 45 Example 6 C 5 820 5 30 120 450 45 Comparative Example 7 C 5 820 90 30 30 450 45 Comparative Example 8 D 735 30 780 150 70 150 450 60 Example 9 D 30 780 120 3 210 450 60 Comparative Example 10 D 30 780 120 100 380 450 50 Comparative Example 11 E 708 7 850 75 80 180 400 30 Exam ple 12 E 7 850 60 80 200 250 60 Comparative Example 13 E 7 830 75 80 200 650 60 Comparative Example 14 E 7 850 75 80 40 400 30 Comparative Example 15 F 723 15 800 240 90 200 400 90 Example 16 F 15 820 240 90 220 400 0 Comparative Example 17 F 15 800 240 90 240 500 900 Comparative Example 18 G 725 15 850 60 100 200 500 30 Example 19 H 720 15 840 120 90 180 400 30 Example 20 I 718 15 830 75 150 220 500 45 Example 21 J 743 15 800 45 80 180 400 20 Example 22 K 730 15 800 200 100 210 550 10 Comparative Example 23 L 686 15 820 120 150 220 400 60 Comparative Example 24 M 745 15 840 90 150 160 400 20 Comparative Example Note: Underlined items are outside the range of the present invention. [Table 3] No. Steel type Ferrite Martensite Tempered martensite Retained austenite Average grain diameter of low-temperature transformation- Other phases TS EL TS×EL Hole expansion ratio Absorption energy up to 10% (AE) AE/TS area% a rea% area% volume% µm MPa % MPa•% % MJ/m 1 A 65 0 29 6 2.7 900 26 23400 85 57 0.063 Example 2 A 63 0 35 2 2.8 915 20 18300 92 58 0.063 Comparative Example 3 B 70 0 26 4 2.4 870 26 22620 88 56 0.064 Example 4 B 73 0 8 0 1.9 P 835 21 17535 60 46 0.055 Comparative Example 5 C 55 0 39 6 3.4 990 23 22770 75 57 0.058 Example 6 C 62 0 9 1 2.6 P 935 20 18700 55 49 0.052 Comparative Example 7 C 57 0 42 1 3.5 980 19 18620 92 56 0.057 Comparative Example 8 D 57 0 31 12 1.7 975 26 25350 81 65 0.067 Example 9 D 65 0 25 1 2.1 P 920 20 18400 63 50 0.054 Comparative Example 10 D 58 20 0 14 1.8 B 970 25 24250 32 65 0.067 Comparative Example 11 E 69 5 21 5 3.6 850 26 22100 89 47 0.055 Example 12 E 70 13 15 2 3.6 859 22 18898 74 49 0.057 Comparative Example 13 E 65 0 20 1 3.7 P 823 23 18929 88 40 0.049 Comparative Example 14 E 75 0 24 1 3.5 820 22 18040 105 44 0.054 Comparative Example 15 F 72 0 21 7 2.1 840 27 22680 74 54 0.064 Example 16 F 70 12 17 1 2.0 865 21 18165 62 56 0.065 Comparative Example 17 F 72 0 18 1 2.1 B 796 23 18308 82 50 0.063 Comparative Example 18 G 53 0 37 10 1.8 1015 26 26390 76 72 0.071 Example 19 H 65 0 30 5 2.2 900 25 22500 95 59 0.066 Example 20 I 51 0 42 7 2.8 1068 23 24564 85 68 0.064 Example 21 J 75 0 20 5 2.7 923 24 22152 92 60 0.065 Example 22 K 91 0 8 1 1.8 611 28 17108 73 33 0.054 Comparative Example 23 L 15 0 76 9 2.9 1325 14 18550 75 69 0.052 Comparative Example 24 M 86 0 5 0 2.7 P 562 30 16860 65 31 0.055 Comparative Example Note: Underlined items are outside the range of the present invention.
* : B represents bainite and P represents pearlite. - The cross-sectional microstructure of the steel sheet was observed by exposing the microstructure by using a 3% nital solution (3% nitric acid + ethanol), observing the position 1/4 of the thickness in the depth direction by using a scanning electron microscope, and conducting an image processing of a picture of the microstructure taken to determine the fraction of the ferrite phase (the image processing can be performed by using commercially available image processing software). The area fractions of the martensite and tempered martensite were determined by taking SEM photographs of adequate magnification, e.g., about 1000 to 3000 magnification, depending on the fineness of the microstructure and then determining the quantity by using image processing software. The average grain diameter of the low-temperature transformation-forming phase was determined by dividing the area of the low-temperature transformation-forming phases in the observed area by the number of the low-temperature transformation-forming phases, determining the average area therefrom, and raising the average to the power of 1/2.
- The volume ratio of the retained austenite was determined by polishing the steel sheet to a surface 1/4 in the thickness direction and measuring X-ray diffraction intensity of the 1/4 thickness surface. A MoKα line was used as the incident X ray, the intensity ratios were determined for all combinations of the integrated intensities of peaks of {111}, {200}, {220}, and {311} faces of the retained austenite phase and the {110}, {200}, and {211} faces of the ferrite phase, and the average value was assumed to be the volume fraction of the retained austenite.
- The tensile property was determined by using a JIS No. 5 specimen sampled from the steel sheet in such a manner that the tensile direction was orthogonal to the rolling direction, conducting a tensile test according to JIS Z2241 to measure TS (tensile strength) and EL (elongation), and determining the strength-elongation balance value represented by the product of the strength and elongation (TS × EL).
- The hole expanding ratio λ was measured as an indicator for evaluating the stretch flangeability. The hole expanding ratio λ was determined by conducting a hole expanding test according to the Japan Iron and Steel Federation standard JFST1001 and determining the ratio from the initial diameter (10 mmφ) of the hole upon punching and the diameter of hole at the time the crack at the hole edge penetrated the sheet upon hole expanding.
- The shock absorption property was determined by using a specimen 5 mm in width and 7 mm in length sampled from the steel sheet in a direction orthogonal to the rolling direction, conducting a tensile test at a strain rate of 2000/s, and integrating the stress-true strain curve obtained by the tensile test within the range of 0 to 10% to calculate the absorption energy (refer to Tetsu-to-Hagane, 83 (1997) p. 748).
- The steel sheets of the examples of the present invention have excellent strength, ductility, and stretch flangeability, i.e., TS × EL of 22000 MPa·% or more and λ of 70% or more.
- In contrast, the steel sheets of comparative examples outside the range of the present invention did not achieve excellent strength, ductility, and stretch flangeability unlike the steel sheets of the examples of the present invention since TS × EL was less than 22000 MPa·% and/or λ was less than 70%. Moreover, when the average grain diameter of the low-temperature transformation-forming phase is 3 µm or less, the ratio of the absorption energy to TS (AE/TS) is 0.063 or more, thereby achieving excellent crashworthiness.
- The present invention can contribute to weight reduction and decreasing the fuel consumption of automobiles by providing a high-strength cold rolled steel sheet having excellent formability and crashworthiness.
Claims (1)
- A method for manufacturing a high-strength cold rolled steel sheet having excellent formability and crashworthiness, the method comprising hot-rolling and cold-rolling a slab having a composition comprising, on a mass% basis, C: 0.05 to 0.3%, Si: 0.3 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100%, S: 0.02% or less, Al: 0.010 to 0.5%, optionally at least one element selected from Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.005 to 2.00%, and Cu: 0.005 to 2.00%, optionally one or both of Ti: 0.01 to 0.20% and Nb: 0.01 to 0.20%, optionally B: 0.0002 to 0.005%, optionally one or both of Ca: 0.001 to 0.005% and REM: 0.001 to 0.005%, and balance being iron and unavoidable impurities, to manufacture a cold rolled steel sheet and continuously annealing the cold rolled sheet, wherein, during the continuous annealing, the steel sheet is heated with an average heating rate in the range of 500°C to Ac1 transformation point being 10°C/s or more, held at a temperature of 750°C or more for 10 seconds or more, cooled from 750°C to a temperature range of 150 to 350°C at a cooling rate of 10°C/s or more on average, heated to a temperature of 350 to 600°C, held thereat for 10 to 600 seconds, and cooled to room temperature, wherein the high-strength cold rolled steel sheet microstructure includes 20% or more of ferrite on an area fraction basis, 10 to 60% of tempered martensite on an area fraction basis, 0 to 10% of martensite on an area fraction basis, and 3 to 15% of retained austenite on a volume fraction basis.
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Families Citing this family (48)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
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CN109072381B (en) | 2016-04-14 | 2020-12-15 | 杰富意钢铁株式会社 | High-strength steel sheet and method for producing same |
WO2017222160A1 (en) * | 2016-06-21 | 2017-12-28 | 현대제철 주식회사 | High-strength cold rolled steel sheet having excellent bendability, and manufacturing method therefor |
US11365465B2 (en) | 2016-08-08 | 2022-06-21 | Nippon Steel Corporation | Steel sheet |
US11136644B2 (en) | 2016-08-31 | 2021-10-05 | Jfe Steel Corporation | High-strength cold rolled steel sheet and method for producing the same |
MX2020005485A (en) | 2017-11-29 | 2020-08-27 | Jfe Steel Corp | High-strength cold-rolled steel sheet and method for manufacturing same. |
JP6597938B1 (en) * | 2018-01-31 | 2019-10-30 | Jfeスチール株式会社 | High-strength cold-rolled steel sheet, high-strength plated steel sheet, and methods for producing them |
EP3778975B1 (en) * | 2018-03-30 | 2024-06-26 | JFE Steel Corporation | High-strength steel sheet and production method thereof |
CN112703265A (en) * | 2018-10-04 | 2021-04-23 | 日本制铁株式会社 | Cold rolled steel sheet |
US20220098698A1 (en) * | 2019-01-29 | 2022-03-31 | Jfe Steel Corporation | High-strength steel sheet and method for producing the same |
CN112760554A (en) * | 2019-10-21 | 2021-05-07 | 宝山钢铁股份有限公司 | High-strength steel with excellent ductility and manufacturing method thereof |
CN117660846A (en) * | 2022-08-23 | 2024-03-08 | 宝山钢铁股份有限公司 | 120 kg-level cold-rolled low-alloy annealed dual-phase steel and manufacturing method thereof |
Family Cites Families (14)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JPH0670247B2 (en) | 1988-10-05 | 1994-09-07 | 新日本製鐵株式会社 | Method for producing high strength steel sheet with good formability |
JP3587126B2 (en) * | 1999-04-21 | 2004-11-10 | Jfeスチール株式会社 | High tensile hot-dip galvanized steel sheet excellent in ductility and method for producing the same |
JP4188608B2 (en) * | 2001-02-28 | 2008-11-26 | 株式会社神戸製鋼所 | High-strength steel sheet with excellent workability and method for producing the same |
ATE383452T1 (en) * | 2001-10-04 | 2008-01-15 | Nippon Steel Corp | DRAWABLE HIGH STRENGTH THIN STEEL SHEET HAVING EXCELLENT FORM-FIXING PROPERTIES AND PRODUCTION PROCESS THEREOF |
JP4119758B2 (en) * | 2003-01-16 | 2008-07-16 | 株式会社神戸製鋼所 | High-strength steel sheet excellent in workability and shape freezing property, and its production method |
JP2004256872A (en) | 2003-02-26 | 2004-09-16 | Jfe Steel Kk | High-tensile strength cold-rolled steel sheet superior in elongation and formability for extension flange, and manufacturing method therefor |
ATE526424T1 (en) * | 2003-08-29 | 2011-10-15 | Kobe Steel Ltd | HIGH EXTENSION STRENGTH STEEL SHEET EXCELLENT FOR PROCESSING AND PROCESS FOR PRODUCTION OF THE SAME |
KR100955982B1 (en) * | 2005-03-31 | 2010-05-06 | 가부시키가이샤 고베 세이코쇼 | High-strength cold-rolled steel sheet excellent in coating adhesion, workability and hydrogen embrittlement resistance, and steel component for automobile |
CN101821419B (en) * | 2007-10-25 | 2015-03-18 | 杰富意钢铁株式会社 | High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same |
JP5369663B2 (en) * | 2008-01-31 | 2013-12-18 | Jfeスチール株式会社 | High-strength hot-dip galvanized steel sheet excellent in workability and manufacturing method thereof |
JP4894863B2 (en) * | 2008-02-08 | 2012-03-14 | Jfeスチール株式会社 | High-strength hot-dip galvanized steel sheet excellent in workability and manufacturing method thereof |
JP5463685B2 (en) * | 2009-02-25 | 2014-04-09 | Jfeスチール株式会社 | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same |
TWI510641B (en) * | 2011-12-26 | 2015-12-01 | Jfe Steel Corp | High strength steel sheet and manufacturing method thereof |
US20150203947A1 (en) * | 2012-07-31 | 2015-07-23 | Jfe Steel Corporation | High-strength galvanized steel sheet with excellent formability and shape fixability and method for manufacturing the same |
-
2010
- 2010-08-12 MX MX2013001456A patent/MX2013001456A/en not_active Application Discontinuation
- 2010-08-12 US US13/816,561 patent/US20130133792A1/en not_active Abandoned
- 2010-08-12 EP EP10855912.1A patent/EP2604715B1/en active Active
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- 2010-08-12 CN CN2010800685780A patent/CN103069040A/en active Pending
- 2010-08-12 KR KR1020137003735A patent/KR20130036763A/en not_active Application Discontinuation
Non-Patent Citations (1)
Title |
---|
None * |
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