Classifications

• • C—CHEMISTRY; METALLURGY
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
  • C23C2/06—Zinc or cadmium or alloys based thereon
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
  • B21C—MANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES, PROFILES OR LIKE SEMI-MANUFACTURED PRODUCTS OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
  • B21C47/00—Winding-up, coiling or winding-off metal wire, metal band or other flexible metal material characterised by features relevant to metal processing only
  • B21C47/02—Winding-up or coiling
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  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • 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
  • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
• • C—CHEMISTRY; METALLURGY
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  • 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")
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  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/74—Methods of treatment in inert gas, controlled atmosphere, vacuum or pulverulent material
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  • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0226—Hot rolling
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  • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0236—Cold rolling
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  • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0242—Flattening; Dressing; Flexing
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  • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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  • C21—METALLURGY OF IRON
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  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys 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
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  • 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
  • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
• • 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
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
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  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium 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
  • 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
  • C23C2/0224—Two or more thermal pretreatments
• • C—CHEMISTRY; METALLURGY
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
• • C—CHEMISTRY; METALLURGY
  • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
  • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
  • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
  • C23C2/26—After-treatment
  • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
  • C23C2/29—Cooling or quenching
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  • 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
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/002—Bainite
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  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite
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  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite

Definitions

• the present disclosure relates to a high-strength steel sheet used for automobile structural members, or the like, and more particularly, a steel sheet having excellent formability and fracture resistance and a method for manufacturing the same.
• high-strength automobile materials may be classified into precipitation-strengthened steel, bake-hardened steel, solid solution-strengthened steel, and transformation-strengthened steel.
• DP steel may ensure high-strength as hard martensite is finely and homogeneously dispersed within soft ferrite
• CP steel may include two or three phases of ferrite, martensite, and bainite, and precipitation hardening elements such as Ti and Nb may be added to DP to improve strength
• TRIP steel may include finely and homogeneously dispersed residual austenite, and when processed at room temperature, the residual austenite phase may transform into martensite, thereby ensuring high-strength and high ductility.
• an ultra-high-strength steel sheet with a tensile strength of 980 MPa or more may be increasingly used for a vehicle body structure or as a reinforcing material.
• high-strength steel having excellent yield strength may be employed for structural members such as a member, a seat rail, and a pillar to improve collision resistance of the vehicle body.
• yield strength may need to be improved relative to tensile strength.
• a representative method for achieving this may be a method of using water cooling during continuous annealing. In this method, cracks may be created in the annealing process, the steel may be immersed in water to form martensite, and through a tempering process, a steel sheet having a microstructure of tempered martensite may be manufactured.
• cited document 1 discloses a martensitic steel having a volume ratio of martensite of 80 to 97% by continuously annealing steel including carbon (C) of 0.18% or more, cooling to room temperature, and overaging at a temperature of 120 to 300°C for 1 to 15 minutes.
• C carbon
• the yield ratio may be relatively high, but the shape quality of the coil may deteriorate due to temperature deviation in the width direction and length direction, such that cracks may be created during forming and workability may be reduced.
• Cited document 2 discloses a method for manufacturing a high-strength steel sheet having a complex phase mainly including martensite, in which fine precipitated copper particles having a particle size of 1 to 100 nm are dispersed in the structure to improve workability.
• fine precipitated copper particles having a particle size of 1 to 100 nm are dispersed in the structure to improve workability.
• Cu may be excessively added, 2-5 weight%, such that red-hot embrittlement may occur due to Cu, and manufacturing costs may increase excessively.
• Cited document 3 is a precipitation-strengthened steel sheet having ferrite as a base structure and including 2 to 10 area% of pearlite, and attempts to improve strength through precipitation strengthening and grain refinement by adding carbon and nitride forming elements such as Nb, Ti, and V.
• This steel sheet may have good hole expandability, but may have limitations in increasing tensile strength, may have high yield strength and low ductility, such that cracks may be created during press forming.
• Cited document 4 discloses a method for manufacturing a cold rolled steel sheet simultaneously ensuring high-strength and high ductility by using tempered martensite and also having an excellent plate shape after continuous annealing.
• the carbon content may be high, 0.2% or more, such that weldability may deteriorate and due to the large amount of Si content, denter defects may occur in the furnace.
• An aspect of the present disclosure is to provide a steel sheet suitable for an automobile structural member, and more particularly, a steel sheet having high strength and also excellent formability and fracture resistance, and a method for manufacturing same.
• An aspect of the present disclosure provides a steel sheet comprising, by weight%, carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities,
• a fraction of fresh martensite satisfies relational expression 2 as below: FM TM + B / FM T ⁇ 100 ⁇ 80 % where FM T indicates a total fraction of fresh martensite, and FM TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite.
• Yield strength (YS), post uniform elongation (P-El) and tensile strength (TS) of the steel sheet satisfy relational expression 4 as below: YS ⁇ P ⁇ El / TS ⁇ 3
• An aspect of the present disclosure provides a method of manufacturing a steel sheet comprising heating a steel slab comprising, by weight%, carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities, wherein C, Si, Al, Mn, Cr and Mo satisfy relational expression 1 as below, in a temperature range of 1100-1300°C;
• the method further comprises holding the steel sheet for 30 seconds after the secondary-cooling.
• the method further comprises hot-dip galvanizing the steel sheet in a plating bath at 430-490°C after the reheating and holding.
• the method further comprises performing alloying heat treatment after the hot-dip galvanizing.
• the method further comprises cooling the steel sheet to a temperature of Ms-100°C or lower at a cooling rate of 5°C/s or higher and performing a skin pass of less than 2% after the alloying heat treatment.
• the steel sheet having high tensile strength of 980 MPa or more and excellent formability, processing defects occurring during processing of parts may be prevented, thereby manufacturing various parts having complex shapes. Also, by improving fracture resistance, the steel sheet may contribute to improvement of safety of parts and vehicle.
• the inventors of the present disclosure have conducted in-depth research to provide a high-strength steel sheet with excellent formability and fracture resistance.
• the content of each element may be based on weight, and the ratio of the structure may be based on area.
• the steel sheet may comprise carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities.
• the C content exceeds 0.2%, strength due to the formation of martensite may increase, but the strength difference with ferrite having a low carbon concentration may increase. Due to this strength difference, fracture may easily occur at an interfacial surface between phases during plastic deformation, ductility and work hardening rate may decrease. Also, weldability may be deteriorated, such that welding defects may occur during processing of parts, and liquid metal embrittlement (LME) cracks may occur during welding, which may impair performance of the parts.
• the C content is less than 0.1%, it may be difficult to ensure a target level of strength, and it may be difficult to ensure a predetermined fraction of residual austenite phase required for ductility. It may be preferable for the C content to be 0.10-0.20%. More preferably, the content may be 0.12-0.18%.
• Si may be a ferrite stabilizing element and may facilitate the transformation of ferrite and may promote enrichment of carbon (C) into untransformed austenite, thereby contributing to the formation of martensite.
• Si may have excellent solid-solution strengthening ability, which may be effective in increasing strength of ferrite and reducing the difference in hardness between phases.
• Si may promote the enrichment of C into untransformed austenite, thereby delaying the transformation of martensite during low-temperature rapid cooling, and forming residual austenite necessary for ductility, and thus, Si may be useful for improving ductility of the steel sheet.
• Si may be useful in that Si may ensure strength without decreasing ductility of the steel sheet.
• the content of Si exceeds 1.3%, surface scale defects may occur, which adversely affect the plating surface quality, may impede chemical processability, and may deteriorate weldability, such that welding defects may occur during processing of parts. In particular, LME cracks may occur during welding, which decreases performance of parts.
• the content is less than 0.5%, it may be difficult to ensure a predetermined fraction of residual austenite phase required for ductility, and solid-solution hardenability may deteriorate, such that strength of ferrite may be lowered, and there may be limitations in reduction of the hardness difference between phases, and accordingly, formability may decrease. It may be more preferable that the content of Si be 0.50-1.30%. More preferably, the content may be 0.7-1.2%.
• Aluminum more preferably acid-soluble aluminum (Sol.Al), may be added for grain refinement and deoxidation of steel, and may be a ferrite stabilizing element similarly to Si.
• Al may be useful for distributing carbon in ferrite to austenite to improve hardenability of martensite. Also, Al may effectively suppress precipitation of carbides in bainite during holding in the bainite region during annealing, thereby promoting C enrichment into untransformed austenite, and delaying martensite transformation during low-temperature rapid cooling, and Al may generate a residual austenite phase, thereby improving ductility of the steel sheet.
• the content of Al exceeds 0.5%, inclusions may be excessively formed during the steelmaking continuous casting operation, it may be highly likely that surface defects may occur on the surface of the steel sheet, and manufacturing costs may also increase. Also, weldability may be deteriorated, which may cause welding defects during processing of parts. Accordingly, the Al may be included in 0.5% or less, and 0% may be excluded. It may be preferable that the Al content be 0.50% or less.
• Mn may refine particles without decreasing ductility and may completely precipitate sulfur (S) in the steel as MnS, thereby preventing hot brittleness caused by FeS formation, and Mn may also be effective for strengthening steel. Also, Mn may facilitate the formation of martensite by lowering a critical cooling rate at which the martensite phase is obtained in complex phase steel.
• Mn content When the Mn content is less than 1.9%, it may be difficult to ensure the target strength in the present disclosure.
• the content exceeds 3.0% it may be highly likely that there may be a problem in weldability and hot-rollability, and martensite may be formed excessively, the material may become unstable, Mn-bands may be formed within the structure, such that there may be a high risk that defects such as processing cracks and plate breakage may occur. Also, during the annealing process, Mn oxide may be dissolved on the surface, which may significantly deteriorate the surface quality.
• Mo may be selectively included to delay the transformation of austenite into pearlite and to improve the refinement and strength of ferrite. Mo may improve hardenability of the steel and may control the yield ratio by forming martensite finely at the grain boundary. However, as Mo is an expensive element, the manufacturing cost may increase as the content thereof increases, which may be economically disadvantageous.
• Cr may be added to improve hardenability of the steel and to ensure high strength, and may be important in formation of martensite. Also, Cr may be preferable for the production of complex phase steel having high ductility by minimizing the decrease in elongation as compared to the increase in strength.
• Cr-based carbides such as Cr 23 C 6 may be formed during the hot-rolling process, and some of these carbides may be dissolved during the annealing process and some may remain undissolved, such that the solid-solution C content in martensite may be controlled below an appropriate level after cooling, such that yield point elongation (YP-El) may be prevented, and thus, Cr may be preferable for the production of complex phase steel having a low yield ratio.
• the Cr content exceeds 1.0%, the above-mentioned effect may be saturated, and the hot-rolled strength may increase excessively and the cold-rollability may deteriorate. Also, the Cr-based carbides may be formed excessively and coarsen, the martensite size may become coarse after annealing, causing a decrease in elongation. It may be more preferable that the Cr content be 1.00% or less.
• Phosphorus (P) 0.1% or less
• P may be a substitutional element having a large solid-solution strengthening effect, and may improving the anisotropy in the plane, may not significantly damage formability and may be most preferable element for ensuring strength.
• the content of the P is excessive, the possibility of brittle fracture may greatly increase, slab plate breakage may occur during hot-rolling and the plating surface properties may be degraded. Accordingly, P may be included at 0.1% or less, and 0% may be excluded considering the level inevitably added during the steel manufacturing process.
• S may be impurities inevitably added to steel, and as S may deteriorate ductility and weldability, and thus, it may be preferable to manage the content thereof as low as possible. In particular, since S may cause red-hot brittleness, it may be preferable to control the content to 0.01% or less. However, considering the level inevitably added during the steel manufacturing process, 0% may be excluded.
• B may delay the transformation of austenite into pearlite during the cooling process during annealing, and may be selectively included to ensure hardenability to suppress formation of ferrite and promote the formation of martensite.
• B may be excessively concentrated on the steel surface, such that plating adhesion may deteriorate. Accordingly, B may be included in 0.01% or less. It may be more preferable that the content of B be 0.010% or less.
• titanium (Ti): 0.05% or less and niobium (Nb): 0.05% or less may be further comprised.
• Ti and Nb may be effective for increasing strength of steel and refining grains by forming nano precipitates.
• the elements may combine with carbon and may form fine nano precipitates, and these nano precipitates may strengthen the base structure and may thus reduce the difference in hardness between phases.
• each of Ti and Nb may be included in 0.05% or less. It may be more preferable that the content of Ti and Nb be 0.050% or less.
• a remainder of the present disclosure is iron (Fe).
• Fe iron
• inevitable impurities may be inevitably added from raw materials or an ambient environment, and thus, impurities may not be excluded.
• a person skilled in the art of a general manufacturing process may be aware of the impurities, and thus, the descriptions of the impurities may not be provided in the present disclosure.
• the steel sheet among the alloy compositions described above, it may be effective for C, Si, Al, Mn, Cr and Mo to satisfy relational expression 1 as below. 10 C + Si + Al / Mn + 1.3 Cr + 2.7 Mo ⁇ 0.7
• relational expression 1 When relational expression 1 is not satisfied, it may be difficult to ensure a sufficient residual austenite fraction and fine fresh martensite may not be uniformly distributed, it may not be easy to ensure formability and fracture resistance. It may be preferable that relational expression 1 be 0.70 or less.
• the microstructure may comprise, by area fraction, ferrite: 10-35%, residual austenite: 3-15%, fresh martensite: 20% or less (excluding 0%), and the remainder may include one or more of tempered martensite and bainite.
• the ferrite may be important for ensuring formability and strength. As the fraction of ferrite increases, carbon may be accumulated in untransformed austenite, such that the martensite transformation temperature may be lowered below room temperature, and may thus contribute to ensuring residual austenite at room temperature. Also, since ferrite is easily transformed into a soft phase, ferrite may contribute to ensuring ductility. To this end, ferrite may be effective at 10% or more. When the fraction of the ferrite exceeds 35%, strength may decrease and the desired strength may not be ensured, and when it is less than 10%, the ductility contribution effect of ferrite may not be expected, and the accumulation of carbon in the austenite may be impaired, such that it may be difficult to ensure the desired fraction of residual austenite.
• bainite when bainite is generated or the generated martensite may be tempered, carbon included in bainite or martensite may move to the surrounding untransformed austenite and may accumulate, such that the martensite transformation temperature may be lowered to room temperature or lower, and accordingly, residual austenite may be ensured at room temperature.
• the residual austenite is 3% or more, transformation-induced plasticity may occur during forming, and it may be preferable to ensure ductility of the steel sheet.
• the residual austenite when the residual austenite is excessive, the plated steel sheet may tend to be vulnerable to liquid metal brittleness during spot welding of automobile parts, such that the content may be effective not to exceed 15%.
• Fresh martensite may be effective in ensuring strength, and when the fraction exceeds 20%, fresh martensite may not be finely and evenly distributed around the tempered martensite or bainite structure.
• a large local stress concentration may be received during forming or impact in accordance with high hardness properties, which may cause brittleness or fracture. In other words, formability may be reduced.
• the remainder of the microstructure be at least one of tempered martensite and bainite phases.
• carbon may be concentrated in untransformed austenite, and residual austenite contributing to ductility may be generated in a fraction of 3-15%, and the fraction of fresh martensite may also be controlled to 20% or less. Accordingly, fresh martensite may be dispersed finely and evenly around tempered martensite or bainite structures.
• the fraction is less than 40%, the desired residual austenite fraction may not be ensured, and the fraction of fresh martensite ultimately generated may also exceed 20%, such that it may be difficult to evenly and finely disperse fine fresh martensite around tempered martensite or bainite.
• fresh martensite may be the microstructure having the highest strength among the microstructure phases.
• the local stress concentration applied to fresh martensite may be distributed to neighboring microstructures and may be alleviated. Accordingly, since the generation and coalescence of pores are delayed, excellent formability may be obtained without processing cracks during the forming of a part.
• stress may be locally concentrated in the fresh martensite such that the creation and coalescence of pores may be facilitated. Accordingly, processing cracks may occur during component molding, which may reduce formability.
• the fresh martensite may have a fraction of fresh martensite in which the distance between fresh martensite phases may be three or more times of the grain size of fresh martensite.
• the fraction of fresh martensite in which the distance between fresh martensite phases (L) is three or more times of the grain size (d) of the fresh martensite (L > 3d) is 30% or more, the local stress concentration in the fresh martensite previously described may be prevented and the effect of delaying the creation and coalescence of pores may be maximized, thereby improving formability and impact fracture resistance.
• the fraction is less than 30%, the stress may be locally concentrated in the fresh martensite may be increased, such that the creation and coalescence of pores may be facilitated, thereby deteriorating formability and impact fracture resistance.
• the fraction of the fresh martensite may satisfy relational expression 2. This may indicate that implementing a microstructure in which fine fresh martensite is uniformly dispersed around tempered martensite or bainite may be effective in terms of formability and fracture resistance.
• the fraction of fresh martensite in contact with tempered martensite or bainite mentioned in relational expression 2 exceeds 80% of the total fraction of fresh martensite, the evenly distributed effect of fine fresh martensite may maximize the local stress concentration relief effect and the creation and coalescence delay effect in fresh martensite, thereby improving formability and impact fracture resistance.
• relational expression 3 may indicate that, when fine fresh martensite may be evenly distributed by controlling the microstructure, excellent formability may be obtained despite high yield strength. Also, when the post uniform elongation is excellent, more deformation may have to accompany until the breakage of the final material even when necking occurs. It may be preferable that relational expression 3 be 6.0 or more.
• Relational expression 4 may indicate that, as mentioned above, when fine fresh martensite is evenly distributed by controlling the microstructure, yield strength may be high and the post uniform elongation may also be excellent, such that the generation and coalescence of pores in the structure during a vehicle collision may be delayed, and impact fracture resistance may thus be excellent. It may be preferable that relational expression 4 be 3.0 or more. In summary, when relational expressions 3 and 4 are satisfied, it may indicate that the steel material may be excellent, which may have excellent part formability, and excellent fracture resistance during a vehicle collision although yield strength is high.
• the steel sheet of the present disclosure may be implemented as a cold-rolled steel sheet, and also a hot-dip galvanized steel sheet including a zinc-based plating layer on at least one surface of the cold-rolled steel sheet, or an alloyed hot-dip galvanized steel sheet alloyed with the hot-dip galvanized steel sheet.
• the zinc-based plating layer may be a zinc-plated layer including mainly zinc, or a zinc alloy plating layer including aluminum and/or magnesium in addition to zinc.
• a steel slab satisfying the aforementioned alloy composition and relational expression 1 may be prepared, and the steel slab may be heated. This process may be performed to smoothly perform the subsequent hot-rolling process and to sufficiently obtain properties of the target steel sheet.
• the heating process may be not specifically limited, and normal heating conditions may be sufficient.
• the heating may be performed in a temperature range of 1100-1300°C.
• the heating temperature is less than 1100°C, friction between the steel sheet and the rolling mill may increase, the load applied to the roller during hot rolling may increase rapidly.
• the temperature exceeds 1300°C, the energy cost required for temperature increase may increase, and the amount of surface scale may also increase, which may lead to material loss.
• the heated steel slab may be finished hot-rolled at the Ar3 transformation point or higher and a hot-rolled steel sheet may be manufactured, and in this case, the hot-rolling conditions may not be particularly limited and may be performed at a normal hot-rolling temperature. As a preferable example, the finishing hot-rolling may be performed at a temperature range of 800-1000°C.
• the hot-rolled steel sheet manufactured as above may be winded, and in this case, the winding may be performed at a temperature range of 400-700°C.
• the strength of the hot-rolled steel sheet may be excessively high, which may cause a rolling load during subsequent cold-rolling.
• the cost and time required to cool the hot-rolled steel sheet to the winding temperature may be excessive, which may increase the process cost.
• the temperature exceeds 700°C excessive scale may occur on the surface of the hot-rolled steel sheet, which may cause surface defects and may deteriorate the plating properties.
• cooling may indicate an average cooling rate.
• a hot-rolled steel sheet in which carbides becoming nucleation sites of austenite are finely dispersed may be obtained. That is, during the hot rolling process, fine carbides may be evenly dispersed in the steel, and the carbides may be dissolved during the subsequent annealing, such that an austenite phase may be finely dispersed and formed in the steel, and accordingly, after the annealing is completed, a uniformly dispersed fine martensite phase may be obtained.
• the wound hot-rolled steel sheet may be cold-rolled to manufacture a cold-rolled steel sheet, and in this case, the cold-rolling may be performed with a cold reduction ratio (total reduction ratio) of 30-80%.
• the stored energy in the steel may be increased, thereby acting as a driving force to promote the recrystallization of ferrite in the subsequent annealing process. Accordingly, the effect of lowering the fraction of unrecrystallized ferrite in the steel may be provided.
• the cold-rolling may be performed using a rolling mill consisting of 5 or 6 stands, but an embodiment thereof is not limited thereto.
• the continuous annealing treatment may be performed, for example, in a continuous alloying galvanizing furnace.
• the continuous annealing may be a process for forming ferrite and austenite phases simultaneously with recrystallization and decomposing carbon.
• the continuous annealing treatment may be performed in a temperature range of Ac1+30°C-Ac3+30°C, and more preferably in a temperature range of 800-880°C.
• the temperature during continuous annealing is lower than Ac1+30°C, recrystallization may not be sufficient, and it may be be difficult to form sufficient austenite, and thus, the target-level fractions of tempered martensite, fresh martensite, and bainite after annealing may not be ensured.
• the temperature exceeds Ac3+30°C, the austenite grain size may become coarse, and the fine residual austenite phase may not be evenly formed around the hard phase.
• the productivity may decrease, and the formation of surface concentrated substances may be intensified due to elements decreasing wettability of hot-dip galvanized such as Si, Mn, and B due to high-temperature annealing, the plating surface quality may not be ensured.
• the cooling may be preferably performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) to 450-670°C (the cooling in this case is referred to as primary-cooling), and at an average cooling rate of 5°C/s or higher to 250-500°C (the cooling in this case is referred to as secondary-cooling).
• the present disclosure may be required to ensure the fractions of martensite and bainite in the subsequent secondary-cooling process in order to form 40% or more of one or more of fractions of tempered martensite and bainite as the final structure.
• the primary-cooling may be performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) to 450-670°C.
• the subsequent secondary-cooling when the subsequent secondary-cooling is terminated below Ms (martensite transformation initiation temperature), a relatively large amount of martensite phase may be formed, and to this end, it may be preferable to control the termination temperature of the primary-cooling to be as low as possible. Also, when the subsequent secondary-cooling is terminated in the bainite temperature region, a bainite phase may be formed relatively preferably, and to this end, it may be preferable to control the termination temperature of the primary-cooling to be higher.
• Ms martensite transformation initiation temperature
• Cooling during primary-cooling may be performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) up to 450-700°C, and preferably, when subsequent secondary-cooling is terminated below Ms, it may be preferable to perform the primary-cooling within a temperature range of 450-600°C, and when subsequent secondary-cooling is terminated in the bainite temperature region, it may be preferable to perform the primary-cooling within a temperature range of 550-700°C.
• the termination temperature during primary-cooling is less than 450°C
• the facility cooling the atmosphere gas in the annealing furnace may be heavily loaded, and the cooling rate may increase such that the ferrite phase formed during cooling may not be sufficiently ensured.
• the termination temperature exceeds 700°C, an excessively high cooling speed may be required during subsequent cooling (secondary-cooling), which may be disadvantageous.
• the primary-cooling may be performed at an average cooling rate of 1°C/s or more.
• quenching martensite When the secondary-cooling is performed below Ms, quenching martensite may be formed, and as the temperature decreases, the fraction of quenching martensite may increase, which may induce improvement in the strength of the steel sheet. Also, during the subsequent heat treatment (the reheating process of the present disclosure), martensite may be tempered and may become tempered martensite, the supersaturated carbon in the martensite may be distributed to the surrounding untransformed austenite, such that stability of the residual austenite may increase and ductility may improve.
• the fraction of bainite may be increased.
• carbon may be distributed from the bainite to the surrounding untransformed austenite, stability of the residual austenite may be increased and ductility may be improved.
• the cooling rate during the secondary-cooling may be higher than the cooling rate during the primary-cooling.
• a process of holding 30 seconds or more in the cooled temperature range may be performed.
• the effect of tempering martensite or further increasing the amount of bainite transformation may be obtained.
• the holding time is less than 30 seconds, it may be difficult to expect the above-described effect.
• the microstructure intended in the present disclosure may be formed through the process of reheating and holding the cold-rolled steel sheet having gone through the sequential cooling as described above. Specifically, it may be preferable to perform a process of reheating the secondary-cooled cold-rolled steel sheet to a temperature of 490°C or lower and holding the steel sheet for 20 seconds or more.
• the quenching martensite generated in the previous cooling process may be transformed into tempered martensite, and also bainite transformation may be accompanied.
• the reheating temperature may be performed at 490°C or lower, and more preferably at 470°C or lower.
• the tempering effect of martensite may become excessive, which decrease the strength, such that it may be preferable not to exceed 5 minutes.
• a plated steel sheet may be provided by plating a cold-rolled steel sheet manufactured as the above.
• the alloyed hot-dip galvanized steel sheet may be obtained by performing alloying heat treatment on the hot-dip galvanized steel sheet.
• the alloying heat treatment process conditions may not be particularly limited, and any normal conditions may be used.
• the alloying heat treatment process may be performed at a temperature range of 480-600°C.
• the final cooling and the skin pass process may be performed after the hot-dip galvanizing or the alloying heat treatment.
• the steel sheet having gone through the hot-dip galvanizing or the alloying heat treatment as above may be finally cooled and fresh martensite may be further formed.
• the final cooling may be preferably performed at a cooling rate of 5°C/s or higher to a temperature of Ms or lower.
• the fresh martensite phase may not be ensured to the intended level during the cooling process.
• the upper limit of the cooling rate is not particularly limited, but the cooing may be performed at 50°C/s or lower to form a predetermined fraction of the pressed martensite phase.
• the reduction ratio be less than 2% (excluding 0%).
• the ratio may be preferable in terms of dislocation formation, but side effects such as plate breakage may occur due to facility capacity limitations.
• each cold-rolled steel sheet was continuously annealed under the conditions listed in Table 2, and primary-cooling, secondary-cooling, holding, and reheating were performed. After the secondary-cooling and holding processes were completed, reheating to a temperature of 490°C or lower was performed, and holding at the above temperature for 30 seconds was performed. The holding process after the secondary-cooling was performed for 30 seconds.
• the mechanical properties and microstructure properties of each sample were evaluated, and the results are listed in Table 3 below.
• the tensile test for each test sample was performed in the L direction using the DIN standard and the tensile properties were evaluated.
• the strain to the breakage was defined as total elongation (T-El)
• the strain to the maximum load point was defined as uniform elongation (U-El)
• the strain from the maximum load point to the breakage was defined as post uniform elongation (P-El), and values thereof were listed.
• the base structure was nital-corroded, and the fractions of tempered martensite (TM), bainite (B), ferrite (F), fresh martensite (FM), and residual austenite (A) were measured using FE-SEM, image analyzer, EBSD, and XRD.
• TM tempered martensite
• B bainite
• F ferrite
• FM fresh martensite
• A residual austenite
• relational expressions 1 to 4 are as below. 10 C + Si + Al / Mn + 1.3 Cr + 2.7 Mo ⁇ 0.7
• FM T indicates a total fraction of fresh martensite
• FM TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite
• FIG. 1 is a graph indicating the relationship between relational expression 1 and relational expression 3 for inventive steels 1 to 6 and comparative steels 1 to 5 in the embodiment
• FIG. 2 may be a graph indicating the relationship between relational expression 1 and relational expression 4 for inventive steel 1 to 6 and comparative steel 1 to 5.
• relational expression 1 of relational expression 1 is 0.7 or higher
• relational expressions 3 and 4 were satisfied.
• relational expression 3 and relational expression 4 may need to be satisfied. To this end, it may be effective to ensure the microstructure presented in the present disclosure, and to this end, it may be important to enrich C in untransformed austenite.
• a predetermined amount of C content may need to be included in the tempering composition, and additionally, the formation of carbides may need to be delayed in the bainite transformation region by adding Si and Al. Also, proper ferrite may need to be ensured to sufficiently enrich C in austenite, but when hardenability elements such as Mn, Cr, and Mo are added excessively, the formation of ferrite may be suppressed, such that C may not be sufficiently enriched in austenite.
• relational expression 1 When relational expression 1 is 0.7 or more, the fractions of the target ferrite and residual austenite may be ensured, and accordingly, fresh martensite may also be finely dispersed, thereby increasing uniform elongation and non-uniform ductility, and as a result, relational expressions 3 and 4 may also be satisfied.
• relational expression 1 When the relational expression 1 is less than 0.7, the target residual austenite and ferrite fraction may not be ensured, and as a result, the fresh martensite may not be finely dispersed, such that the uniform elongation and post uniform elongation may decrease, and relational expression 3 and relational expression 4 may not be satisfied.

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