JP5456026B2 - High-strength steel sheet, hot-dip galvanized steel sheet with excellent ductility and no cracks at the edge, and manufacturing method thereof - Google Patents
High-strength steel sheet, hot-dip galvanized steel sheet with excellent ductility and no cracks at the edge, and manufacturing method thereof Download PDFInfo
- Publication number
- JP5456026B2 JP5456026B2 JP2011510432A JP2011510432A JP5456026B2 JP 5456026 B2 JP5456026 B2 JP 5456026B2 JP 2011510432 A JP2011510432 A JP 2011510432A JP 2011510432 A JP2011510432 A JP 2011510432A JP 5456026 B2 JP5456026 B2 JP 5456026B2
- Authority
- JP
- Japan
- Prior art keywords
- steel sheet
- hot
- rolled steel
- strength
- cooling
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Active
Links
- 229910000831 Steel Inorganic materials 0.000 title claims description 97
- 239000010959 steel Substances 0.000 title claims description 97
- 238000004519 manufacturing process Methods 0.000 title claims description 22
- 229910001335 Galvanized steel Inorganic materials 0.000 title claims description 8
- 239000008397 galvanized steel Substances 0.000 title claims description 8
- 238000001816 cooling Methods 0.000 claims description 51
- 229910001566 austenite Inorganic materials 0.000 claims description 37
- 229910000734 martensite Inorganic materials 0.000 claims description 35
- 238000000034 method Methods 0.000 claims description 25
- 238000005097 cold rolling Methods 0.000 claims description 24
- 238000000137 annealing Methods 0.000 claims description 23
- 238000005098 hot rolling Methods 0.000 claims description 22
- 229910000859 α-Fe Inorganic materials 0.000 claims description 21
- 230000009467 reduction Effects 0.000 claims description 16
- 229910052782 aluminium Inorganic materials 0.000 claims description 14
- 239000010960 cold rolled steel Substances 0.000 claims description 14
- 229910052710 silicon Inorganic materials 0.000 claims description 13
- 229910052748 manganese Inorganic materials 0.000 claims description 12
- 239000012535 impurity Substances 0.000 claims description 10
- 229910052796 boron Inorganic materials 0.000 claims description 9
- 229910052719 titanium Inorganic materials 0.000 claims description 9
- 238000004804 winding Methods 0.000 claims description 7
- 229910001563 bainite Inorganic materials 0.000 claims description 6
- 238000005096 rolling process Methods 0.000 claims description 4
- 238000005246 galvanizing Methods 0.000 claims description 2
- 230000015572 biosynthetic process Effects 0.000 description 16
- 230000008569 process Effects 0.000 description 15
- 230000000717 retained effect Effects 0.000 description 14
- 239000011572 manganese Substances 0.000 description 13
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 12
- 229910052799 carbon Inorganic materials 0.000 description 12
- 230000000694 effects Effects 0.000 description 10
- 230000009466 transformation Effects 0.000 description 10
- 239000000463 material Substances 0.000 description 8
- 150000001247 metal acetylides Chemical class 0.000 description 8
- 239000000203 mixture Substances 0.000 description 8
- 239000010936 titanium Substances 0.000 description 8
- 238000007747 plating Methods 0.000 description 7
- 230000007423 decrease Effects 0.000 description 6
- 229910001562 pearlite Inorganic materials 0.000 description 5
- 238000012360 testing method Methods 0.000 description 5
- 239000006104 solid solution Substances 0.000 description 4
- 230000000052 comparative effect Effects 0.000 description 3
- 239000010410 layer Substances 0.000 description 3
- 238000005728 strengthening Methods 0.000 description 3
- 229910000794 TRIP steel Inorganic materials 0.000 description 2
- 229910000797 Ultra-high-strength steel Inorganic materials 0.000 description 2
- 238000005275 alloying Methods 0.000 description 2
- 229910052787 antimony Inorganic materials 0.000 description 2
- -1 carbides Chemical class 0.000 description 2
- 238000007796 conventional method Methods 0.000 description 2
- 238000010438 heat treatment Methods 0.000 description 2
- 230000000087 stabilizing effect Effects 0.000 description 2
- ZOXJGFHDIHLPTG-UHFFFAOYSA-N Boron Chemical compound [B] ZOXJGFHDIHLPTG-UHFFFAOYSA-N 0.000 description 1
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 1
- PWHULOQIROXLJO-UHFFFAOYSA-N Manganese Chemical compound [Mn] PWHULOQIROXLJO-UHFFFAOYSA-N 0.000 description 1
- XUIMIQQOPSSXEZ-UHFFFAOYSA-N Silicon Chemical compound [Si] XUIMIQQOPSSXEZ-UHFFFAOYSA-N 0.000 description 1
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 description 1
- HCHKCACWOHOZIP-UHFFFAOYSA-N Zinc Chemical compound [Zn] HCHKCACWOHOZIP-UHFFFAOYSA-N 0.000 description 1
- XAGFODPZIPBFFR-UHFFFAOYSA-N aluminium Chemical compound [Al] XAGFODPZIPBFFR-UHFFFAOYSA-N 0.000 description 1
- WATWJIUSRGPENY-UHFFFAOYSA-N antimony atom Chemical compound [Sb] WATWJIUSRGPENY-UHFFFAOYSA-N 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
- 238000004090 dissolution Methods 0.000 description 1
- 238000009826 distribution Methods 0.000 description 1
- 230000009977 dual effect Effects 0.000 description 1
- 229910001039 duplex stainless steel Inorganic materials 0.000 description 1
- 238000003912 environmental pollution Methods 0.000 description 1
- 238000005530 etching Methods 0.000 description 1
- 239000000446 fuel Substances 0.000 description 1
- 239000007789 gas Substances 0.000 description 1
- 229910052739 hydrogen Inorganic materials 0.000 description 1
- 239000001257 hydrogen Substances 0.000 description 1
- 239000004615 ingredient Substances 0.000 description 1
- 239000011159 matrix material Substances 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 238000012986 modification Methods 0.000 description 1
- 230000004048 modification Effects 0.000 description 1
- 230000003287 optical effect Effects 0.000 description 1
- 238000001953 recrystallisation Methods 0.000 description 1
- 230000002040 relaxant effect Effects 0.000 description 1
- 238000011160 research Methods 0.000 description 1
- 238000005204 segregation Methods 0.000 description 1
- 239000010703 silicon Substances 0.000 description 1
- 238000010583 slow cooling Methods 0.000 description 1
- 239000000243 solution Substances 0.000 description 1
- 239000000126 substance Substances 0.000 description 1
- 239000002344 surface layer Substances 0.000 description 1
- 239000013077 target material Substances 0.000 description 1
- 238000009864 tensile test Methods 0.000 description 1
- 230000001131 transforming effect Effects 0.000 description 1
- 239000011701 zinc Substances 0.000 description 1
- 229910052725 zinc Inorganic materials 0.000 description 1
Images
Classifications
-
- 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
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- 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
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0236—Cold rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0273—Final recrystallisation annealing
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/04—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
- C21D8/0421—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 characterised by the working steps
- C21D8/0436—Cold rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/04—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
- C21D8/0447—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 characterised by the heat treatment
- C21D8/0463—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 characterised by the heat treatment following hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/04—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
- C21D8/0447—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 characterised by the heat treatment
- C21D8/0473—Final recrystallisation annealing
-
- 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
- C22C38/14—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
- C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
-
- 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
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
- C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
- C23C2/0224—Two or more thermal pretreatments
-
- C—CHEMISTRY; METALLURGY
- 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
-
- 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
-
- Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
- Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
- Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
- Y10T428/00—Stock material or miscellaneous articles
- Y10T428/12—All metal or with adjacent metals
- Y10T428/12493—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, joint, etc.]
- Y10T428/12771—Transition metal-base component
- Y10T428/12785—Group IIB metal-base component
- Y10T428/12792—Zn-base component
- Y10T428/12799—Next to Fe-base component [e.g., galvanized]
Landscapes
- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Chemical Kinetics & Catalysis (AREA)
- Heat Treatment Of Sheet Steel (AREA)
- Coating With Molten Metal (AREA)
Description
本発明は、主に自動車の各種構造部材に用いられる強度及び加工性に優れた自動車用鋼板及びその製造方法に関し、より詳しくは、780〜980MPaの引張強度と少なくとも28%以上の伸び率とを有し、かつ冷間圧延時に圧延板のエッジ(Edge)部位に亀裂が発生しない、高強度冷延鋼板及び溶融亜鉛メッキ鋼板を製造する方法に関する。 The present invention relates to a steel sheet for automobiles excellent in strength and workability mainly used in various structural members of automobiles, and more specifically, a tensile strength of 780 to 980 MPa and an elongation of at least 28% or more. The present invention relates to a method for producing a high-strength cold-rolled steel sheet and a hot-dip galvanized steel sheet that does not crack at the edge (Edge) portion of the rolled sheet during cold rolling.
自動車の構成部品において構造部材に用いられる鋼板は、外部からの衝撃をうまく吸収することによって車両衝突時に乗客の安全性を向上させる役割をする。最近、このような部品としては、例えば引張強度780MPa以上で、かつ伸び率が25%以上のような、強度・成形性ともに優れた鋼板が主に用いられている。このような鋼板は、高引張強度と高伸び率、低降伏比(降伏強度/引張強度の比率)を有する必要がある。 Steel plates used as structural members in automobile components play a role in improving passenger safety in the event of a vehicle collision by well absorbing external impacts. Recently, as such parts, steel sheets having excellent strength and formability such as tensile strength of 780 MPa or more and elongation rate of 25% or more are mainly used. Such a steel plate needs to have high tensile strength, high elongation, and low yield ratio (yield strength / tensile strength ratio).
さらに、近年、自動車排気ガスによる環境汚染問題が注目を浴びており、より燃量消費を低減させるために、引張強度780MPa以上の超高強度鋼を用いた、自動車の軽量化のための研究が増加しつつある。 Furthermore, in recent years, environmental pollution problems due to automobile exhaust gas have attracted attention, and in order to further reduce fuel consumption, research for reducing the weight of automobiles using ultra-high strength steel with a tensile strength of 780 MPa or more has been conducted. It is increasing.
自動車用高強度鋼は、一般的に変態誘起塑性(TRIP:Transformation Induced Plasticity)鋼と、二相(DP:Dual Phase)鋼とを含む。 High-strength steel for automobiles generally includes transformation induced plasticity (TRIP) steel and dual phase (DP) steel.
加工性に優れた超高強度鋼の製造工程は、スラブ製造工程、熱間圧延工程、熱間圧延板材を冷却して巻取する工程、冷間圧延工程、そして焼鈍工程に大別される。熱間圧延したフェライトとパーライト組織を有する板材を冷間圧延して加工し、冷間圧延した板材をA1変態点〜A3変態点間の温度で焼鈍した後、冷却時の鋼板の冷却速度を調節して、焼鈍工程において形成されたオーステナイトをマルテンサイトに変態させて得られる鋼を、二相鋼という。上記二相鋼の強度は、二相鋼中に含まれるマルテンサイトとフェライトの分率により決定される。全体構造中の、マルテンサイトの比率が増加するほど、強度が増加し延性は減少する。このため、二相鋼は適正なマルテンサイトの比率を有しなければならない。一方、上記の二相鋼を製造する方法のように、焼鈍過程においてオーステナイトを形成後、冷却工程において冷却速度と終了冷却温度を制御することにより常温でオーステナイトの一部を残留させることによって、鋼材の強度と延性を同時に増加させる方法がある。このように生成された残留オーステナイトは、鋼材の強度を増加させるため塑性変形中にマルテンサイトに変態させるのと同時に、塑性誘起変態により応力集中を緩和させることで延性を増加させる。これが高強度と高延性を有する高強度鋼として用いられている、変態誘起塑性(TRIP)鋼である。 The manufacturing process of ultra-high-strength steel excellent in workability is roughly divided into a slab manufacturing process, a hot rolling process, a process of cooling and winding a hot-rolled plate material, a cold rolling process, and an annealing process. A plate having a hot rolled ferrite and pearlite processed by cold rolling, after the cold rolled sheet was annealed at a temperature between A 1 transformation point to A 3 transformation point, the cooling rate of the steel plate during cooling The steel obtained by transforming the austenite formed in the annealing process to martensite by adjusting the temperature is called a duplex steel. The strength of the duplex steel is determined by the fraction of martensite and ferrite contained in the duplex steel. As the proportion of martensite in the overall structure increases, the strength increases and the ductility decreases. For this reason, the duplex stainless steel must have the proper martensite ratio. On the other hand, by forming austenite in the annealing process as in the above-described method for producing the duplex steel, by controlling the cooling rate and the end cooling temperature in the cooling process, a part of the austenite remains at room temperature, thereby making the steel material There is a method of simultaneously increasing the strength and ductility of the steel. Residual austenite generated in this manner is transformed into martensite during plastic deformation in order to increase the strength of the steel material, and at the same time, ductility is increased by relaxing stress concentration by plastic-induced transformation. This is a transformation induced plasticity (TRIP) steel used as a high strength steel having high strength and high ductility.
上記の変態誘起塑性鋼の強度と延性を高めるには、常温において準安定性残留オーステナイトをマルテンサイトに変態するのを抑制して、一定の分率以上の残留オーステナイトを維持することが非常に重要である。このため、Si、Alなどを添加して、フェライト内の炭素の活性(activity)を高めて炭化物の生成を抑制し、オーステナイト相内の炭素の濃度を増加させることによって、残留オーステナイトの安定性を確保する。 In order to increase the strength and ductility of the above transformation-induced plastic steel, it is very important to suppress metastable retained austenite to martensite at room temperature and maintain retained austenite above a certain fraction It is. Therefore, the stability of retained austenite is increased by adding Si, Al, etc. to increase the activity of carbon in the ferrite to suppress the formation of carbides and increase the concentration of carbon in the austenite phase. Secure.
このような高強度と伸び率とを確保するための既知の技術は、特許文献1に開示の技術を含む。この技術によると、C:0.6〜0.22重量(wt)%、Si:0.05〜1.0wt%、Mn:0.5〜2.0wt%、及びAl:0.25〜1.5wt%を含む鋼に、Mo:0.03〜0.3wt%を添加して、490MPa以上の強度と35%以上の伸び率とを有する鋼が提供される。 Known techniques for ensuring such high strength and elongation include the technique disclosed in Patent Document 1. According to this technique, C: 0.6 to 0.22 wt (wt)%, Si: 0.05 to 1.0 wt%, Mn: 0.5 to 2.0 wt%, and Al: 0.25 to 1 By adding Mo: 0.03 to 0.3 wt% to steel containing 0.5 wt%, a steel having a strength of 490 MPa or more and an elongation of 35% or more is provided.
また、特許文献2では、C:0.15〜0.30wt%、Si:1.5〜2.5wt%、Mn:0.5〜2.0wt%、及びAl:0.02〜0.1wt%を含み、かつ780MPa以上の強度と30%以上の伸び率とを有する鋼が開示されている。
In
特許文献3では、C:0.06〜0.6wt%、[Si+Al]:0.5〜3.0wt%、Mn:0.5〜3.0wt%を含み、かつ約800MPaの強度と40%の伸び率とを有する鋼が開示されている。しかし、特許文献3に開示されている成分の実際の内容を見ると、[Al+Si]の量は1.5wt%以下であることが分かる。例えば、Alが1wt%添加されると、Siは0.5wt%添加される。従って、[Si+Al]の成分範囲は0.5〜3.0wt%ではなく、1.5wt%が正確な表現である。 In Patent Document 3, C: 0.06 to 0.6 wt%, [Si + Al]: 0.5 to 3.0 wt%, Mn: 0.5 to 3.0 wt%, and a strength of about 800 MPa and 40% Steels having an elongation of about 5 are disclosed. However, when the actual content of the component currently disclosed by patent document 3 is seen, it turns out that the quantity of [Al + Si] is 1.5 wt% or less. For example, when 1 wt% of Al is added, 0.5 wt% of Si is added. Therefore, the component range of [Si + Al] is not 0.5 to 3.0 wt%, but 1.5 wt% is an accurate expression.
また、特許文献3に開示されている製造工程は二つに分けられるが、その一つは熱間圧延後に焼鈍やメッキを行う製造工程であり、もう一つは、熱間圧延、冷間圧延後、2段階焼鈍を行う製造工程である。2段階焼鈍において、冷間圧延後の第1段階の焼鈍によってマトリックス組織内にマルテンサイトが形成され、第2段階の焼鈍は通常の変態誘起塑性鋼で施す焼鈍を行う。これは経済的な損失をもたらし、かつ適用可能性が低い。 Moreover, although the manufacturing process currently disclosed by patent document 3 is divided into two, one is a manufacturing process which performs annealing and plating after hot rolling, and the other is hot rolling and cold rolling. Thereafter, it is a manufacturing process for performing two-stage annealing. In the two-stage annealing, martensite is formed in the matrix structure by the first-stage annealing after the cold rolling, and the second-stage annealing is performed by an ordinary transformation-induced plastic steel. This results in economic losses and is less applicable.
上記した従来技術は、強度がそれぞれ約490MPa〜780MPa程度に過ぎず、かつ製造工程が複雑である。したがって、780MPa以上の引張強度と28%以上の伸び率とを同時に確保するためには、経済的に有利で適用可能性も高い製造技術が必要とされている。 The above-described conventional techniques each have a strength of only about 490 MPa to 780 MPa, and the manufacturing process is complicated. Therefore, in order to simultaneously secure a tensile strength of 780 MPa or more and an elongation of 28% or more, a manufacturing technique that is economically advantageous and highly applicable is required.
本発明の目的は、組成成分を制御し、熱間圧延での冷却段階を調整してマルテンサイト分率を30〜70%含むように制御することで、780〜980MPaの引張強度と28%以上の伸び率とを有すると共に、エッジ部位に亀裂が発生しない高強度鋼板、溶融亜鉛メッキ鋼板及びその製造方法を提供することにある。 The object of the present invention is to control the composition components and adjust the cooling stage in hot rolling so as to include a martensite fraction of 30 to 70%, thereby providing a tensile strength of 780 to 980 MPa and 28% or more. It is to provide a high-strength steel plate, a hot-dip galvanized steel plate, and a method for producing the same, in which an edge portion has no cracks.
本発明の一態様では、重量%で、C:0.1〜0.25%、Si:1.0〜1.9%、Mn:1.5〜2.5%、Al:0.5〜1.6%、Ti:0.005〜0.03%、B:5〜30ppm、及びSb:0.01〜0.03%、並びに残部Fe及び不可避の不純物からなり、かつ1.75≦Si+Al≦3.25%を満たす高強度熱延鋼板であって、熱間圧延後の前記鋼板の微細組織が、30〜70%のマルテンサイトと、残部フェライトからなることを特徴とする、高強度鋼板及び溶融亜鉛メッキ鋼板を提供する。 In one embodiment of the present invention, C: 0.1 to 0.25%, Si: 1.0 to 1.9%, Mn: 1.5 to 2.5%, Al: 0.5 to 1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and inevitable impurities, and 1.75 ≦ Si + Al A high-strength hot-rolled steel sheet satisfying ≦ 3.25%, wherein the microstructure of the steel sheet after hot rolling is composed of 30-70% martensite and the balance ferrite. And a galvanized steel sheet.
本発明の別の態様では、重量%で、C:0.1〜0.25%、Si:1.0〜1.9%、Mn:1.5〜2.5%、Al:0.5〜1.6%、Ti:0.005〜0.03%、B:5〜30ppm、及びSb:0.01〜0.03%、並びに残部Fe及び不可避の不純物からなり、1.75≦Si+Al≦3.25%を満たす鋼スラブを、A3以上の温度で熱間圧延する段階と、30〜200℃/s(秒)の冷却速度で600〜800℃の温度範囲で1次冷却する段階と、上記1次冷却した熱延鋼板を600〜800℃の温度範囲で空冷する段階と、上記空冷した熱延鋼板を50〜200℃/sの冷却速度で常温〜300℃の温度範囲で2次冷却する段階と、上記2次冷却した熱延鋼板を常温〜300℃の温度範囲で巻取する段階と、を含むことを特徴とする、高強度熱延鋼板を製造する方法が提供される。 In another embodiment of the present invention, by weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5 ~1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, 1.75 ≦ Si + Al the steel slab satisfying ≦ 3.25%, comprising the steps of hot rolling at a 3 or higher, the step of primary cooling in the temperature range of 600 to 800 ° C. at a cooling rate of 30 to 200 ° C. / s (sec) And air cooling the primary-cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C, and the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C / s at a temperature range of room temperature to 300 ° C. And a step of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C. A method for producing a high-strength hot-rolled steel sheet is provided.
また、本発明の別の態様では、重量%で、C:0.1〜0.25%、Si:1.0〜1.9%、Mn:1.5〜2.5%、Al:0.5〜1.6%、Ti:0.005〜0.03%、B:5〜30ppm、及びSb:0.01〜0.03%、並びに残部Fe及び不可避の不純物からなり、かつ1.75≦Si+Al≦3.25%を満たす鋼スラブを、A3以上の温度範囲で熱間圧延する段階と、30〜200℃/sの冷却速度で600〜800℃の温度範囲で1次冷却する段階と、上記1次冷却した熱延鋼板を600〜800℃の温度範囲で空冷する段階と、上記空冷した熱延鋼板を50〜200℃/sの冷却速度で常温〜300℃の温度範囲で2次冷却する段階と、上記2次冷却した熱延鋼板を常温〜300℃の温度範囲で巻取する段階と、上記巻取した熱延鋼板を30〜50%の圧下率で冷間圧延する段階と、前記冷間圧延した鋼板を焼鈍する段階と、を含むことを特徴とする高強度冷延鋼板を製造する方法が提供される。 In another embodiment of the present invention, by weight, C: 0.1 to 0.25%, Si: 1.0 to 1.9%, Mn: 1.5 to 2.5%, Al: 0 .5~1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, and 1. A steel slab satisfying 75 ≦ Si + Al ≦ 3.25% is subjected to primary rolling in a temperature range of 600 to 800 ° C. at a stage of hot rolling in a temperature range of A 3 or higher and a cooling rate of 30 to 200 ° C./s. Air cooling the primary-cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C., and air-cooling the hot-rolled steel sheet in a temperature range of room temperature to 300 ° C. at a cooling rate of 50 to 200 ° C./s. A stage of secondary cooling, a stage of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C .; Manufacturing a high-strength cold-rolled steel sheet, comprising: cold-rolling the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%; and annealing the cold-rolled steel sheet. A method is provided.
本発明の別の態様では、上記の高強度冷延鋼板を製造する方法に加えて、溶融亜鉛メッキする段階をさらに含む、高強度溶融亜鉛メッキ鋼板を製造する方法が提供される。 In another aspect of the present invention, there is provided a method for producing a high-strength hot-dip galvanized steel sheet, further comprising the step of hot-dip galvanizing in addition to the above-described method for producing a high-strength cold-rolled steel sheet.
本発明の上記の、並びに他の態様、特徴、及び他の利点は、添付の図面と共に、以下の発明の詳細な説明によって明確に理解されるだろう。 The above as well as other aspects, features, and other advantages of the present invention will become apparent from the following detailed description of the invention when taken in conjunction with the accompanying drawings.
本発明の実施の形態を図面を参照して詳細に説明する。しかし、本発明は、様々な形で具体化することができ、かつここに記載の実施形態に限定されるものと解釈すべきではない。むしろ、これらの実施形態は、本発明の開示が、徹底的かつ完全であること、及び当業者に本発明の範囲を十分に伝えるように提供されている。図面において、層及び領域の厚さはわかりやすくするために誇張されている。図面の参照番号が成分を表すので、これらの説明は省略される。 Embodiments of the present invention will be described in detail with reference to the drawings. However, the present invention can be embodied in various forms and should not be construed as limited to the embodiments set forth herein. Rather, these embodiments are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the invention to those skilled in the art. In the drawings, the thickness of layers and regions are exaggerated for clarity. Since the reference numerals in the drawings represent the components, these descriptions are omitted.
以下、本発明の組成範囲について詳細に説明する(以下、重量(wt)%)。 Hereinafter, the composition range of the present invention will be described in detail (hereinafter, weight (wt)%).
炭素(C)の含量は0.1〜0.25%とする。Cは、最も重要な成分であり、強度及び延性などのあらゆる物理的・化学的特性と密接な関係を有する。本発明の鋼板では、炭素の量が0.1%未満であると残留オーステナイトの分率と安定性が減少し、0.25%を超過すると溶接性が低下し、第2相分率の過大な増加により加工性が低下する短所を有する。したがって、Cの組成範囲は0.1〜0.25%に制限する。 The content of carbon (C) is 0.1 to 0.25%. C is the most important component and has a close relationship with all physical and chemical properties such as strength and ductility. In the steel sheet of the present invention, if the amount of carbon is less than 0.1%, the fraction and stability of retained austenite decrease, and if it exceeds 0.25%, the weldability decreases and the second phase fraction is excessive. However, there is a disadvantage that the workability is reduced due to a large increase. Therefore, the composition range of C is limited to 0.1 to 0.25%.
シリコン(Si)の含量は1.0〜1.9%とする。Siはフェライトに固溶されてフェライトを安定化させる成分である。本発明の冷延鋼板では、フェライト内に固溶されて炭素の活性を高めることで、オーステナイト相内の炭素の濃度を増加させて、カーバイド相の生成を抑制させて残留オーステナイトの安定性を高める役割をする。また、固溶により鋼板の強度が増加する効果もある。Siの含量が1.0%未満であると鋼の強度が低下し、カーバイド相のような炭化物の生成を抑制する効果が低下し、1.9%を超過すると熱延スケールを誘発させ、メッキ性が悪くなり、また溶接性も劣化される。したがって、Siの組成範囲は1.0〜1.9%に制限する。 The content of silicon (Si) is 1.0 to 1.9%. Si is a component that is dissolved in ferrite and stabilizes the ferrite. In the cold-rolled steel sheet of the present invention, the concentration of carbon in the austenite phase is increased by increasing the activity of carbon by being dissolved in ferrite, thereby suppressing the formation of carbide phase and increasing the stability of retained austenite. Play a role. Moreover, there exists an effect which the intensity | strength of a steel plate increases by solid solution. If the Si content is less than 1.0%, the strength of the steel is reduced, and the effect of suppressing the formation of carbides such as carbide phase is reduced. If it exceeds 1.9%, a hot rolling scale is induced and plating is performed. The weldability is deteriorated and the weldability is also deteriorated. Therefore, the composition range of Si is limited to 1.0 to 1.9%.
マンガン(Mn)の含量は1.5〜2.5%とする。Mnは硬化性を大きくして針状フェライト及びベイナイトのような低温変態相を容易に生成させ、強度を増加させる成分である。Mnはまた、オーステナイトも安定化させる。したがって、Mnを1.5%未満で添加する場合は上記効果が期待できず、2.5%を超過して添加すると溶接性が低下し、熱間圧延時に板材の中央に偏析帯を形成し、かつ介在物の形成により水素脆性を引き起こす。したがって、Mnの含量は1.5〜2.5%に制限する。 The content of manganese (Mn) is 1.5 to 2.5%. Mn is a component that increases the strength by increasing the curability and easily generating low-temperature transformation phases such as acicular ferrite and bainite. Mn also stabilizes austenite. Therefore, when Mn is added at less than 1.5%, the above effect cannot be expected. When Mn is added at more than 2.5%, the weldability is lowered, and a segregation zone is formed in the center of the plate during hot rolling. In addition, the formation of inclusions causes hydrogen embrittlement. Therefore, the Mn content is limited to 1.5 to 2.5%.
アルミニウム(Al)の含量は0.5〜1.6%とする。AlはSiより固溶強化の効果が弱いが、フェライト安定化元素として固溶強化の効果を示すフェライト安定化元素であり、カーバイドのような炭化物の生成を抑制し、残留オーステナイト中の炭素濃度を増加させて、安定性を高める成分である。Alの含量が0.5%未満であるとオーステナイトの安定性が低下するため、炭化物の生成を抑制し難くなり、また1.6%を超過するとオーステナイトの分率が低下するため、鋼板の延性が相対的に低下し、かつ表面特性が悪くなる。したがって、Alの含量は0.5〜1.6%に制限する。 The content of aluminum (Al) is 0.5 to 1.6%. Al is less effective in solid solution strengthening than Si, but as a ferrite stabilizing element, it is a ferrite stabilizing element that shows the effect of solid solution strengthening. It suppresses the formation of carbides such as carbides, and reduces the carbon concentration in residual austenite. It is a component that increases the stability and increases the stability. If the Al content is less than 0.5%, the stability of austenite decreases, so it becomes difficult to suppress the formation of carbides. If it exceeds 1.6%, the austenite fraction decreases, so the ductility of the steel sheet. Is relatively lowered and the surface properties are deteriorated. Therefore, the Al content is limited to 0.5 to 1.6%.
チタニウム(Ti)の含量は0.005〜0.03%とする。Tiは、AlとNとの結合によって生じるAlN窒化物の形成を抑制し、Al本来の作用ができるようにTiNを形成する成分である。Tiが0.005%未満であるとそのような役割が期待できず、0.03%を超過して添加しても、それ以上の効果は期待できない。したがって、Tiの含量は0.005〜0.03%に制限する。 The content of titanium (Ti) is 0.005 to 0.03%. Ti is a component that forms TiN so as to suppress the formation of AlN nitride caused by the bonding of Al and N and to allow Al to function. If Ti is less than 0.005%, such a role cannot be expected, and even if added in excess of 0.03%, no further effect can be expected. Therefore, the Ti content is limited to 0.005 to 0.03%.
ホウ素(B)の含量は5〜30ppmとする。Bは鋼中に少量を添加しても硬化性を向上させる成分であり、5ppm以上添加されると、高温でオーステナイト粒界に偏析することで、フェライト形成を抑制して硬化性を向上させるが、30ppmを超えて添加すると、鋼板の再結晶温度を上昇させるため、鋼板の絞り性を低下し、かつ溶接性を劣化させる。したがって、Bの含量は5〜30ppmに制限する。 The content of boron (B) is 5 to 30 ppm. B is a component that improves the curability even if a small amount is added to the steel. When 5 ppm or more is added, it segregates at the austenite grain boundary at a high temperature, thereby suppressing ferrite formation and improving the curability. When added over 30 ppm, the recrystallization temperature of the steel sheet is increased, so that the drawability of the steel sheet is lowered and the weldability is deteriorated. Therefore, the B content is limited to 5-30 ppm.
アンチモン(Sb)の含量は0.01〜0.03%とする。Sbは適正量である0.01〜0.03%を添加すると、鋼板の表面特性を改善させるが、0.03%を超えて添加すると、鋼板の表面にSb濃縮が発生し表面特性が悪くなる。Sbは鉄鋼材料に不可避的に含まれる不純物の一種であり、0.01%未満のSb含量が特定の目的なしに製造した鋼中に見出される。Sbの含量が0.01%未満であると表面に濃縮は発生するが、表面特性の変化を起こすには極めて少ない量であるため、0.01%以上とする。したがって、Sbの含量は0.01〜0.03%に制限する。 The content of antimony (Sb) is 0.01 to 0.03%. When Sb is added in an appropriate amount of 0.01 to 0.03%, the surface properties of the steel sheet are improved. However, when Sb is added in an amount exceeding 0.03%, Sb concentration occurs on the surface of the steel plate, resulting in poor surface properties. Become. Sb is a kind of impurity inevitably contained in steel materials, and an Sb content of less than 0.01% is found in steel produced without a specific purpose. When the Sb content is less than 0.01%, concentration occurs on the surface. However, since the amount is extremely small to cause changes in surface characteristics, the content is made 0.01% or more. Therefore, the Sb content is limited to 0.01 to 0.03%.
本発明では1.75≪Si+Al≪3.25を満たす。SiとAlはいずれも鋼の炭化物形成を抑制する役割をし、残留オーステナイト内の固溶炭素の含量を高めて、残留オーステナイトの安定性を改善させる役割をする。よって、TRIP鋼の強度と伸び率とを考慮する時には、これら二つの成分の含量を同時に制御する必要があるが、2成分の含量の合計が3.25%を超過すると、表面酸化物の形成によりメッキ性などの表面品質が低下し、2相域焼鈍熱処理中に起きるオーステナイト分率の減少により強度及び延性が減少することがある。また、2成分の含量の合計が1.75%未満であると、目標とする引張強度780MPa以上のTRIP鋼を製造するのに必要な基本固溶強化の効果が減少し、残留オーステナイトの安定性を確保し難くなる。したがって、Si+Alの含量は1.75〜3.25%に制限する。 In the present invention, 1.75 << Si + Al << 3.25 is satisfied. Both Si and Al play a role in suppressing the carbide formation of the steel, and increase the content of solute carbon in the retained austenite, thereby improving the stability of the retained austenite. Therefore, when considering the strength and elongation rate of TRIP steel, it is necessary to control the contents of these two components at the same time. However, if the total content of the two components exceeds 3.25%, the formation of surface oxides As a result, the surface quality such as plating property is deteriorated, and the strength and ductility may be reduced due to the reduction of the austenite fraction occurring during the two-phase annealing treatment. Further, if the total content of the two components is less than 1.75%, the effect of basic solid solution strengthening necessary for producing a TRIP steel having a target tensile strength of 780 MPa or more is reduced, and the stability of retained austenite is reduced. It becomes difficult to secure. Therefore, the content of Si + Al is limited to 1.75 to 3.25%.
本発明は、上記組成、並びに残部Fe及び不可避の不純物を含む。 The present invention includes the above composition, the balance Fe and inevitable impurities.
以下、本発明の製造方法について詳しく説明する。 Hereinafter, the production method of the present invention will be described in detail.
本発明は、上記組成を満たす鋼スラブをA3以上のオーステナイト領域で熱間圧延後、30〜200℃/sの冷却速度で1次冷却する。このとき、冷却速度が30℃/s未満であるとパーライト組織が形成され、目標とする材質を確保し難く、200℃/sを超過すると鋼板に温度偏差が発生し、残留応力によって素材の歪みが発生することがある。このため、冷却速度は30〜200℃/秒に制限する。 The present invention, a steel slab satisfying the above composition after hot rolling at A 3 above the austenite region, to primary cooling at a cooling rate of 30 to 200 ° C. / s. At this time, when the cooling rate is less than 30 ° C./s, a pearlite structure is formed, and it is difficult to secure a target material. When the cooling rate exceeds 200 ° C./s, a temperature deviation occurs in the steel sheet, and the material is distorted by residual stress. May occur. For this reason, a cooling rate is restrict | limited to 30-200 degreeC / sec.
上記1次冷却後、鋼板は600〜800℃の温度範囲で維持する。これは、鋼板が強制冷却することなく常温の大気で自然対流により冷却されることを意味する。温度が600℃未満であるとフェライト相形成分率が確保し難く、800℃を超過すると過度なフェライトが形成され、またはパーライト組織が形成されることがある。 After the primary cooling, the steel sheet is maintained in a temperature range of 600 to 800 ° C. This means that the steel sheet is cooled by natural convection in a normal temperature atmosphere without forced cooling. If the temperature is less than 600 ° C., it is difficult to secure the ferrite phase formation fraction, and if it exceeds 800 ° C., excessive ferrite may be formed or a pearlite structure may be formed.
上記空冷後、鋼板は50〜200℃/sの冷却速度で2次冷却した後、常温〜300℃で巻取する。冷却速度が50℃/s未満であるとベイナイト相の形成により目標とする組織を得難く、200℃/sを超過すると過度なマルテンサイト相が形成され、圧延板の形状歪みが発生することがある。したがって、2次冷却速度は、50〜200℃/sに制限する。 After the air cooling, the steel sheet is secondarily cooled at a cooling rate of 50 to 200 ° C./s, and then wound at room temperature to 300 ° C. When the cooling rate is less than 50 ° C./s, it is difficult to obtain a target structure due to the formation of the bainite phase, and when it exceeds 200 ° C./s, an excessive martensite phase is formed, and the shape distortion of the rolled sheet may occur. is there. Therefore, the secondary cooling rate is limited to 50 to 200 ° C./s.
また、巻取温度が300℃を超過すると、ベイナイト相の形成により目標とする組織を確保し難く、常温〜300℃で巻取すると基地組織がラス(lath)形状の微細なマルテンサイト組織を有し、熱延後の高転位密度と均一な固溶炭素分布を有するようになる。したがって、巻取温度は常温〜300℃に制限する。 In addition, when the coiling temperature exceeds 300 ° C, it is difficult to secure a target structure due to the formation of a bainite phase, and when coiling at room temperature to 300 ° C, the base structure has a fine lath-shaped martensite structure. However, it has a high dislocation density after hot rolling and a uniform solute carbon distribution. Therefore, the coiling temperature is limited to room temperature to 300 ° C.
巻取後は30〜50%の圧下率(reduction rate)で冷間圧延を行う。通常は60%の圧下率で冷間圧延を行うが、本発明では30〜50%の圧下率で行う。これにより、冷間圧延によって巻取後に生成されたマルテンサイト組織とフェライト組織のうち、フェライト組織に転位密度を十分に増加させることができ、焼鈍中に炭素の再固溶とオーステナイト相の形成が均一に発生する。冷間圧下率が30%未満であるとフェライト組織内の転位密度が十分でないため、目標とする組織が得られず、また50%を超過するとマルテンサイト相とフェライト相の境界で微細な亀裂が発生しやすく、特に圧延板エッジ部位に亀裂欠陥が発生する。したがって、冷間圧下率は30〜50%に制限する。 After winding, cold rolling is performed at a reduction rate of 30 to 50%. Usually, cold rolling is performed at a reduction rate of 60%, but in the present invention, it is performed at a reduction rate of 30 to 50%. As a result, among the martensite structure and ferrite structure generated after coiling by cold rolling, the dislocation density can be sufficiently increased in the ferrite structure, and the re-dissolution of carbon and the formation of the austenite phase during annealing It occurs uniformly. If the cold rolling reduction is less than 30%, the dislocation density in the ferrite structure is not sufficient, so that the target structure cannot be obtained, and if it exceeds 50%, fine cracks occur at the boundary between the martensite phase and the ferrite phase. It tends to occur, and crack defects occur particularly at the rolled plate edge. Therefore, the cold reduction rate is limited to 30 to 50%.
上記の冷間圧延後は、通常の焼鈍工程を行う。 After the cold rolling, a normal annealing process is performed.
本発明では、冷間圧延後に溶融亜鉛メッキまたは合金化溶融亜鉛メッキを行う。焼鈍工程後は溶融亜鉛が入ったメッキ浴に鋼板を通過させて、表面層にメッキ層を一定の厚さで付着させる。このとき、メッキ浴の温度は450〜500℃が好ましく、30℃/s以下の速度で徐冷する工程により、溶融亜鉛メッキ鋼板を製造する。 In the present invention, hot dip galvanization or alloying hot dip galvanization is performed after cold rolling. After the annealing step, the steel sheet is passed through a plating bath containing molten zinc, and the plating layer is adhered to the surface layer with a certain thickness. At this time, the temperature of the plating bath is preferably 450 to 500 ° C., and a hot-dip galvanized steel sheet is produced by a step of slow cooling at a rate of 30 ° C./s or less.
また、メッキ浴槽を通過したメッキ鋼板を直に500〜600℃の温度範囲で加熱して溶融亜鉛メッキ層の合金化熱処理を行った後、30℃/s以下の速度で徐冷させる工程により、合金化溶融亜鉛メッキ鋼板を製造する。 In addition, after the plated steel sheet that has passed through the plating bath is directly heated in a temperature range of 500 to 600 ° C. to perform alloying heat treatment of the hot dip galvanized layer, it is gradually cooled at a rate of 30 ° C./s or less, Produces galvannealed steel sheet.
以下、本方法により製造される鋼板の組織について詳しく説明する。 Hereinafter, the structure of the steel sheet produced by this method will be described in detail.
本発明において熱間圧延後の熱延鋼板はマルテンサイト分率が30〜70%であることを特徴とする。熱間圧延後、1次及び2次冷却工程で生成されたラス形状の微細なマルテンサイト組織は冷間圧延後、焼鈍工程において均一なオーステナイト変態を誘導し、これにより安定化された残留オーステナイトの分率を増加させる。従来の方法では1回の冷却過程だけで熱延鋼板を製造するが、この場合熱延鋼板は粗大な炭化物が存在するパーライト組織を含有するようになる。粗大な炭化物は冷間圧延後の焼鈍過程において再固溶されるが、炭化物の大きさが粗大なため、700℃以上の高温でも再固溶されずに残留しやすく、焼鈍中にオーステナイトは、主に炭素濃度の高い炭化物付近に形成され始め、相対的に炭化物のない領域にはオーステナイトが形成され難くなる。したがって、2相域焼鈍中にオーステナイト分率が低くなり局所的な偏差も発生し、焼鈍後の冷却工程で発生する残留オーステナイト分率が減少するようになる。熱延鋼板でのマルテンサイト相はこのような短所を改善できる。 In the present invention, the hot-rolled steel sheet after hot rolling has a martensite fraction of 30 to 70%. After hot rolling, the lath-shaped fine martensite structure produced in the primary and secondary cooling processes induces uniform austenite transformation in the annealing process after cold rolling, and thereby stabilized residual austenite Increase the fraction. In the conventional method, a hot-rolled steel sheet is produced by only one cooling process. In this case, the hot-rolled steel sheet contains a pearlite structure in which coarse carbides are present. Coarse carbide is re-dissolved in the annealing process after cold rolling, but because the size of the carbide is coarse, it remains easily without being re-dissolved even at a high temperature of 700 ° C. or higher. Austenite is hardly formed in a region where there is relatively no carbide, starting to form mainly in the vicinity of a carbide having a high carbon concentration. Accordingly, the austenite fraction becomes lower during the two-phase annealing and a local deviation occurs, and the residual austenite fraction generated in the cooling step after annealing decreases. The martensite phase in the hot-rolled steel sheet can improve such disadvantages.
本発明では、1次及び2次冷却工程を経てマルテンサイト相を形成させると、炭化物がほとんど形成されない状態で熱延鋼板を製造できる。このようなマルテンサイト相を含んだ熱延鋼板は、焼鈍熱処理中に転位密度の高いマルテンサイト組織で微細な炭化物が形成されると直に再固溶される。このため、本発明の鋼板中の炭素濃度の偏差は、パーライト組織を有する鋼板に比べてかなり減少する。したがって、焼鈍中に形成されるオーステナイト相は結晶粒界付近やラスマルテンサイト組織であった付近に均一に分布して形成されるため、残留オーステナイト分率と安定性を高めるこができる。 In the present invention, when the martensite phase is formed through the primary and secondary cooling steps, the hot-rolled steel sheet can be produced in a state where carbides are hardly formed. Such a hot-rolled steel sheet containing a martensite phase is re-dissolved immediately when fine carbides are formed in a martensite structure having a high dislocation density during annealing heat treatment. For this reason, the deviation of the carbon concentration in the steel sheet of the present invention is considerably reduced as compared with the steel sheet having a pearlite structure. Therefore, since the austenite phase formed during annealing is uniformly distributed and formed in the vicinity of the grain boundaries and the lath martensite structure, the retained austenite fraction and stability can be improved.
また、マルテンサイト相により高くなった転位密度を活用して、比較的低い冷間圧下率を適用することで、冷間圧延後エッジ部位に発生しやすい微細な亀裂形成を抑制できる。このとき、熱延鋼板のマルテンサイト分率が30%未満であると残留オーステナイト分率が増加する効果が微弱であり、また70%以上であると冷間圧延においてエッジ部位に微細な亀裂が発生する。したがって、本発明では熱延鋼板のマルテンサイト組織分率を30〜70%とする。 In addition, by utilizing a relatively low cold reduction rate by utilizing the dislocation density that is higher due to the martensite phase, it is possible to suppress the formation of fine cracks that are likely to occur at the edge portion after cold rolling. At this time, if the martensite fraction of the hot-rolled steel sheet is less than 30%, the effect of increasing the retained austenite fraction is weak, and if it is 70% or more, fine cracks are generated at the edge portion in cold rolling. To do. Therefore, in the present invention, the martensite structure fraction of the hot-rolled steel sheet is set to 30 to 70%.
本発明の冷間圧延及び焼鈍工程を経た冷延鋼板は、フェライトとベイナイト組織に残留オーステナイトが生成される。このとき本発明の鋼板中の残留オーステナイトの分率は5〜15%である。残留オーステナイト相は、熱延鋼板でのラス形状のマルテンサイト組織の影響によりラス形状に現われ、他の形状の残留オーステナイトに比べて安定したものとなる。また、ベイナイト組織の分率は20〜40%であり、残りはフェライトからなる。 In the cold-rolled steel sheet that has undergone the cold rolling and annealing processes of the present invention, retained austenite is generated in the ferrite and bainite structure. At this time, the fraction of retained austenite in the steel sheet of the present invention is 5 to 15%. The retained austenite phase appears in the lath shape due to the influence of the lath-shaped martensite structure in the hot-rolled steel sheet, and is more stable than the remaining austenite in other shapes. Moreover, the fraction of a bainite structure is 20 to 40%, and the remainder consists of ferrite.
本発明の方法によると、加工性に優れた780〜980MPaの引張強度と28%以上の伸び率とを有する高強度冷延鋼板を製造できる。 According to the method of the present invention, a high-strength cold-rolled steel sheet having excellent workability and a tensile strength of 780 to 980 MPa and an elongation of 28% or more can be produced.
以下、本発明を、実施例を参照してより具体的に説明する。 Hereinafter, the present invention will be described more specifically with reference to examples.
(実施例)
表1は、実施例で使用された鋼の組成範囲である。本発明では目標とする引張強度と伸び率とを得るために、成分の適切な調整が重要である。
(Example)
Table 1 shows the composition range of the steel used in the examples. In the present invention, in order to obtain the target tensile strength and elongation, appropriate adjustment of the components is important.
表1の組成を有する鋼スラブを、熱延工程を経て冷却及び巻取後、冷間圧延、焼鈍を行い、これに対する冷却条件と冷間圧下率及び引張試験結果などを、表2に示した。 The steel slab having the composition shown in Table 1 is cooled and wound through a hot rolling process, and then cold-rolled and annealed. The cooling conditions, the cold reduction rate, the tensile test results, and the like are shown in Table 2. .
材料の機械的特性における引張特性は、成分と微細組織の製造条件によって変化する。本発明の鋼板は熱延直後に生成されたラス形状の微細なマルテンサイト組織により焼鈍過程で均一なオーステナイト変態を誘導し、これにより安定化された残留オーステナイトの分率を増加させて、変形を受けたときに変態して優れた機械的特性を得るものである。 The tensile properties in the mechanical properties of the material vary depending on the ingredients and the microstructure production conditions. The steel sheet of the present invention induces uniform austenite transformation in the annealing process due to the fine martensitic structure of lath formed immediately after hot rolling, thereby increasing the fraction of stabilized retained austenite and causing deformation. When transformed, it transforms to obtain excellent mechanical properties.
また、マルテンサイト相により高くなった転位密度を活用して比較的低い冷間圧下率を適用することで、冷間圧延後エッジ部に発生しやすい微細な亀裂形成を抑制しようとした。このような特性を得るためには、基本的に本発明で提案した成分を満たさなければならず、熱延直後の冷却過程とマルテンサイト分率及び冷間圧下率を最適に調率しなければならない。 In addition, by utilizing a relatively low cold reduction rate by utilizing the dislocation density that has become higher due to the martensite phase, an attempt was made to suppress the formation of fine cracks that are likely to occur at the edge after cold rolling. In order to obtain such characteristics, the components proposed in the present invention must be basically satisfied, and the cooling process immediately after hot rolling, the martensite fraction, and the cold reduction rate must be adjusted optimally. Don't be.
表2に、熱延直後の冷却過程とマルテンサイト分率、冷間圧下率、及び機械的性質とエッジ部位の亀裂の発生程度を測定し、その結果を示した。 Table 2 shows the results of the measurement of the cooling process immediately after hot rolling, the martensite fraction, the cold reduction ratio, the mechanical properties, and the degree of occurrence of cracks at the edge portion.
表2に示すように、比較材9から15は、基本的な成分要件を満たさず、製造過程で冷却条件とマルテンサイト分率及び冷間圧下率の条件が適切に調節されていても、本発明において目標とする強度と延性を得ることはできなかったことが分かる。 As shown in Table 2, the comparative materials 9 to 15 do not satisfy the basic component requirements, and even if the cooling conditions, the martensite fraction, and the cold rolling reduction conditions are appropriately adjusted during the manufacturing process, It can be seen that the target strength and ductility could not be obtained in the invention.
また、比較材である2−2、2−3、3−2、3−3、4−2、4−3、5−2、7−2、7−3、8−2は成分要件を満たしているが、冷却条件、マルテンサイト分率、冷間圧下率などの要件には満たさず、かつ目標とする機械的性質が得られず、またはエッジ部位に微細な亀裂が発生したことが分かる。 Moreover, 2-2, 2-3, 3-2, 3-3, 4-2, 4-3, 5-2, 7-2, 7-3, and 8-2 which are comparative materials satisfy the component requirements. However, it can be seen that the requirements such as the cooling condition, the martensite fraction, the cold reduction rate, etc. are not satisfied, and the target mechanical properties cannot be obtained, or fine cracks are generated at the edge portion.
表2においてエッジ部位の亀裂長さの測定は、マルテンサイト分率を200倍率の光学顕微鏡により観察したものである。 In Table 2, the crack length at the edge portion was measured by observing the martensite fraction with an optical microscope of 200 magnification.
特に、厚さのt/4位置での微細組織を2%のナイタール(Nital)エッチング液でエッチングした後、画像分析器で測定した結果である。エッジ部位に発生した微細な亀裂は、冷間圧延された圧延板エッジ部位をランダムに選定して長さ100mm内で発生した最も長い亀裂を30個以上選択した後、平均値とした。 In particular, this is a result obtained by measuring a fine structure at a thickness of t / 4 with a 2% Nital etching solution and then measuring with an image analyzer. The fine cracks generated in the edge part were selected as an average value after randomly selecting the cold-rolled rolled sheet edge part and selecting 30 or more longest cracks generated within a length of 100 mm.
本発明の鋼板は、成分構成と製造条件を制御することによって、780〜980MPaの引張強度と28%以上の伸び率とを有する。従って、本発明の鋼板は、高い強度と加工性を有する構造用部品に用いることができる。また、本発明の鋼板は、エッジ部位に亀裂が発生することを防止することで経済性を高める効果がある。 The steel sheet of the present invention has a tensile strength of 780 to 980 MPa and an elongation of 28% or more by controlling the component structure and production conditions. Therefore, the steel sheet of the present invention can be used for structural parts having high strength and workability. Moreover, the steel plate of this invention has an effect which improves economical efficiency by preventing that a crack generate | occur | produces in an edge part.
本発明は、例示的な実施の形態と関連して示され、かつ記載されているが、特許請求の範囲によって定義される発明の精神および範囲から逸脱することなく、変更および変形ができることは、当業者に明らかであろう。 While the invention has been shown and described in connection with exemplary embodiments, it will be understood that changes and modifications can be made without departing from the spirit and scope of the invention as defined by the claims. It will be apparent to those skilled in the art.
Claims (7)
1次冷却した熱延鋼板を600〜800℃の温度範囲で空冷する段階と、
空冷した熱延鋼板を50〜200℃/sの冷却速度で常温〜300℃の温度範囲で2次冷却する段階と、
2次冷却した熱延鋼板を常温〜300℃の温度範囲で巻取する段階
を含むことを特徴とする、請求項1記載の高強度熱延鋼板の製造方法。 By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005 to 0.03%, B: 5 to 30 ppm, and Sb: 0.01 to 0.03%, balance Fe and inevitable impurities, and satisfy 1.75 ≦ Si + Al ≦ 3.25% the steel slab comprising the steps of hot rolling at a 3 or more temperature and primary cooling in a temperature range of 600 to 800 ° C. at a cooling rate of 30 to 200 ° C. / s,
Air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C .;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
The method for producing a high-strength hot-rolled steel sheet according to claim 1, comprising a step of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C.
1次冷却した熱延鋼板を、600〜800℃の温度範囲で空冷する段階と、
空冷した熱延鋼板を、50〜200℃/sの冷却速度で常温〜300℃の温度範囲で2次冷却する段階と、
2次冷却した熱延鋼板を、常温〜300℃の温度範囲で巻取する段階と、
巻取した熱延鋼板を、30〜50%の圧下率で冷間圧延し焼鈍する段階
とを含むことを特徴とする、請求項2又は3記載の高強度冷延鋼板の製造方法。 By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005 to 0.03%, B: 5 to 30 ppm, and Sb: 0.01 to 0.03%, balance Fe and inevitable impurities, and satisfy 1.75 ≦ Si + Al ≦ 3.25% the steel slab was hot rolled at a 3 or higher, at a cooling rate of 30 to 200 ° C. / s, the method comprising the primary cooling in the temperature range of 600 to 800 ° C.,
A step of air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
Winding the secondary-cooled hot-rolled steel sheet in a temperature range from room temperature to 300 ° C .;
The method for producing a high-strength cold-rolled steel sheet according to claim 2, comprising cold rolling and annealing the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%.
1次冷却した熱延鋼板を600〜800℃の温度範囲で空冷する段階と、
空冷した熱延鋼板を、50〜200℃/sの冷却速度で常温〜300℃の温度範囲で2次冷却する段階と、
2次冷却した熱延鋼板を、常温〜300℃の温度範囲で巻取する段階と、
巻取した熱延鋼板を、30〜50%の圧下率で冷間圧延し焼鈍する段階と、
焼鈍した冷延鋼板を溶融亜鉛メッキする段階
とを含むことを特徴とする、請求項4記載の高強度溶融亜鉛メッキ鋼板の製造方法。 By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005-0.03%, B: 5-30 ppm, and Sb: 0.01-0.03%, balance Fe and inevitable impurities, and satisfying 1.75 ≦ Si + Al ≦ 3.25% the slab was hot rolled at a 3 or higher, at a cooling rate of 30 to 200 ° C. / s, the method comprising the primary cooling in the temperature range of 600 to 800 ° C.,
Air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C .;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
Winding the secondary-cooled hot-rolled steel sheet in a temperature range from room temperature to 300 ° C .;
Cold rolling and annealing the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%,
A method for producing a high-strength hot-dip galvanized steel sheet according to claim 4, comprising the step of hot-dip galvanizing the annealed cold-rolled steel sheet.
Applications Claiming Priority (3)
Application Number | Priority Date | Filing Date | Title |
---|---|---|---|
KR10-2008-0050366 | 2008-05-29 | ||
KR1020080050366A KR101008099B1 (en) | 2008-05-29 | 2008-05-29 | High strength steel sheet amd galvenized steel sheet having excellent ducility and free edge crack and method of manufacturing the same |
PCT/KR2009/002810 WO2009145562A2 (en) | 2008-05-29 | 2009-05-27 | High-strength steel sheet with excellent ductility and crackless edge portion, hot-dip galvanized steel sheet, and manufacturing method thereof |
Publications (2)
Publication Number | Publication Date |
---|---|
JP2011523443A JP2011523443A (en) | 2011-08-11 |
JP5456026B2 true JP5456026B2 (en) | 2014-03-26 |
Family
ID=41377784
Family Applications (1)
Application Number | Title | Priority Date | Filing Date |
---|---|---|---|
JP2011510432A Active JP5456026B2 (en) | 2008-05-29 | 2009-05-27 | High-strength steel sheet, hot-dip galvanized steel sheet with excellent ductility and no cracks at the edge, and manufacturing method thereof |
Country Status (4)
Country | Link |
---|---|
US (1) | US20110064968A1 (en) |
JP (1) | JP5456026B2 (en) |
KR (1) | KR101008099B1 (en) |
WO (1) | WO2009145562A2 (en) |
Families Citing this family (7)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
KR101290426B1 (en) * | 2011-06-28 | 2013-07-26 | 현대제철 주식회사 | High strength hot-rolled steel sheet and method of manufacturing the hot-rolled steel sheet |
KR101359281B1 (en) * | 2011-12-20 | 2014-02-06 | 주식회사 포스코 | Steel sheet having excellent spot weldabity, strength and ductility for automobile and method for manufacturing the same |
JP6056745B2 (en) * | 2013-12-12 | 2017-01-11 | Jfeスチール株式会社 | High formability and high strength cold-rolled steel sheet excellent in chemical conversion treatment and production method thereof |
KR101657784B1 (en) * | 2014-11-28 | 2016-09-20 | 주식회사 포스코 | High ductility and strength cold rolled steel sheet with reduced cracking in hot-rolling and method for manufacturing the same |
KR101630976B1 (en) * | 2014-12-08 | 2016-06-16 | 주식회사 포스코 | Ultra-high strenth galvanized steel sheet having excellent surface and coating adheision and method for manufacturing thereof |
KR102046544B1 (en) * | 2015-04-15 | 2019-11-19 | 닛폰세이테츠 가부시키가이샤 | Hot rolled steel sheet and manufacturing method |
CN114630914A (en) * | 2019-08-19 | 2022-06-14 | 美国钢铁公司 | High strength steel product and annealing method for producing the same |
Family Cites Families (11)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
KR100481366B1 (en) * | 2000-12-08 | 2005-04-07 | 주식회사 포스코 | A hot rolled trip steel sheet with excellent ductility, and a method for manufacturing it |
CN1204284C (en) * | 2000-12-29 | 2005-06-01 | 新日本制铁株式会社 | High-strength molten-zinc-plated steel plate excellent in deposit adhesion and suitability for press forming and process for producing the same |
JP3809074B2 (en) * | 2001-03-30 | 2006-08-16 | 新日本製鐵株式会社 | High-strength hot-dip galvanized steel sheet with excellent plating adhesion and press formability 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 |
CA2387322C (en) * | 2001-06-06 | 2008-09-30 | Kawasaki Steel Corporation | High-ductility steel sheet excellent in press formability and strain age hardenability, and method for manufacturing the same |
US6733874B2 (en) * | 2001-08-31 | 2004-05-11 | Mitsubishi Materials Corporation | Surface-coated carbide alloy cutting tool |
JP2005336526A (en) | 2004-05-25 | 2005-12-08 | Kobe Steel Ltd | High strength steel sheet having excellent workability and its production method |
JP4640130B2 (en) * | 2005-11-21 | 2011-03-02 | Jfeスチール株式会社 | High-strength cold-rolled steel sheet with small variation in mechanical properties and method for producing the same |
EP1960562B1 (en) * | 2005-12-09 | 2015-08-26 | Posco | High strenght cold rolled steel sheet having excellent formability and coating property, zinc-based metal plated steel sheet made of it and the method for manufacturing thereof |
KR100711468B1 (en) | 2005-12-23 | 2007-04-24 | 주식회사 포스코 | High strength cold rolled steel sheet and hot dip galvanized steel sheet having excellent formability and coating property, and the method for manufacturing thereof |
KR100931140B1 (en) * | 2006-10-31 | 2009-12-10 | 현대자동차주식회사 | High tensile steel sheet with excellent formability and manufacturing method thereof |
-
2008
- 2008-05-29 KR KR1020080050366A patent/KR101008099B1/en active IP Right Grant
-
2009
- 2009-05-27 JP JP2011510432A patent/JP5456026B2/en active Active
- 2009-05-27 US US12/991,003 patent/US20110064968A1/en not_active Abandoned
- 2009-05-27 WO PCT/KR2009/002810 patent/WO2009145562A2/en active Application Filing
Also Published As
Publication number | Publication date |
---|---|
KR20090124264A (en) | 2009-12-03 |
WO2009145562A3 (en) | 2010-01-21 |
WO2009145562A2 (en) | 2009-12-03 |
KR101008099B1 (en) | 2011-01-13 |
JP2011523443A (en) | 2011-08-11 |
US20110064968A1 (en) | 2011-03-17 |
Similar Documents
Publication | Publication Date | Title |
---|---|---|
KR101399741B1 (en) | High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same | |
JP5042232B2 (en) | High-strength cold-rolled steel sheet excellent in formability and plating characteristics, galvanized steel sheet using the same, and method for producing the same | |
KR101638719B1 (en) | Galvanized steel sheet and method for manufacturing the same | |
JP5157215B2 (en) | High rigidity and high strength steel plate with excellent workability | |
KR101264574B1 (en) | Method for producing high-strength steel plate having superior deep drawing characteristics | |
JP5217395B2 (en) | High strength cold-rolled steel sheet with small in-plane anisotropy of elongation and method for producing the same | |
JP5564432B2 (en) | High-strength cold-rolled steel sheet excellent in workability, galvanized steel sheet, and manufacturing method thereof | |
JP4772497B2 (en) | High-strength cold-rolled thin steel sheet excellent in hole expansibility and manufacturing method thereof | |
JP2006104532A (en) | High strength thin steel sheet having excellent elongation and hole expansibility and method for producing the same | |
JP5456026B2 (en) | High-strength steel sheet, hot-dip galvanized steel sheet with excellent ductility and no cracks at the edge, and manufacturing method thereof | |
JP5256690B2 (en) | High-strength hot-dip galvanized steel sheet excellent in workability and impact resistance and method for producing the same | |
KR20140044938A (en) | High strength hot dip galvanized steel sheet having excellent deep-drawability, and method for producing same | |
JP2007177271A (en) | High-strength cold-rolled steel sheet superior in hole expandability, and manufacturing method therefor | |
JP6221424B2 (en) | Cold rolled steel sheet and method for producing the same | |
JP5397141B2 (en) | Alloyed hot-dip galvanized steel sheet and method for producing the same | |
JP4265152B2 (en) | High-tensile cold-rolled steel sheet with excellent elongation and stretch flangeability and method for producing the same | |
CN116507753A (en) | Ultra-high strength steel sheet having excellent ductility and method for manufacturing same | |
JP2010100896A (en) | High strength cold rolled steel sheet having excellent stability in mechanical property and method of producing the same | |
JP4337604B2 (en) | Strain aging treatment method for high-tensile steel sheet and method for producing high-strength structural member | |
JP4010131B2 (en) | Composite structure type high-tensile cold-rolled steel sheet excellent in deep drawability and manufacturing method thereof | |
JP2005290485A (en) | Strain aging treatment method for steel plate and method for manufacturing high-strength structural member | |
JP5304522B2 (en) | High-strength steel sheet with excellent workability and method for producing the same | |
JP5251206B2 (en) | High-strength steel sheet excellent in deep drawability, aging resistance and bake hardenability, and its manufacturing method | |
RU2788613C1 (en) | Cold-rolled coated steel sheet and method for production thereof | |
JP2005232483A (en) | High-strength steel sheet having superior deep drawability and balance of strength and ductility, and manufacturing method therefor |
Legal Events
Date | Code | Title | Description |
---|---|---|---|
A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20101122 |
|
A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20121219 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20130315 |
|
A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20130625 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20130919 |
|
A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20131105 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20131118 |
|
TRDD | Decision of grant or rejection written | ||
A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20131210 |
|
A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20140107 |
|
R150 | Certificate of patent or registration of utility model |
Free format text: JAPANESE INTERMEDIATE CODE: R150 Ref document number: 5456026 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
S531 | Written request for registration of change of domicile |
Free format text: JAPANESE INTERMEDIATE CODE: R313531 |
|
S533 | Written request for registration of change of name |
Free format text: JAPANESE INTERMEDIATE CODE: R313533 |
|
R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
S111 | Request for change of ownership or part of ownership |
Free format text: JAPANESE INTERMEDIATE CODE: R313113 |
|
R371 | Transfer withdrawn |
Free format text: JAPANESE INTERMEDIATE CODE: R371 |
|
S111 | Request for change of ownership or part of ownership |
Free format text: JAPANESE INTERMEDIATE CODE: R313113 |
|
R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |