WO2016045264A1 - 一种高成形性的冷轧超高强度钢板、钢带及其制造方法 - Google Patents

一种高成形性的冷轧超高强度钢板、钢带及其制造方法 Download PDF

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WO2016045264A1
WO2016045264A1 PCT/CN2015/070666 CN2015070666W WO2016045264A1 WO 2016045264 A1 WO2016045264 A1 WO 2016045264A1 CN 2015070666 W CN2015070666 W CN 2015070666W WO 2016045264 A1 WO2016045264 A1 WO 2016045264A1
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cold
steel strip
rolled
steel sheet
strength
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French (fr)
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朱晓东
李伟
薛鹏
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Baoshan Iron and Steel Co Ltd
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Baoshan Iron and Steel Co Ltd
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Priority to MX2017003993A priority Critical patent/MX2017003993A/es
Priority to KR1020177007995A priority patent/KR20170063613A/ko
Priority to US15/514,509 priority patent/US20170298466A1/en
Publication of WO2016045264A1 publication Critical patent/WO2016045264A1/zh
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • 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
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • 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 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
    • 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 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
    • 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 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
    • 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 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
    • 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 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
    • 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 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
    • C21D8/04—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 to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0447—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 to produce plates or strips for drawing, e.g. 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • 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
    • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
    • Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
    • Y02P10/00—Technologies related to metal processing
    • Y02P10/20—Recycling

Definitions

  • the invention relates to a high-formability cold-rolled ultra-high-strength steel plate, a steel strip and a manufacturing method thereof.
  • the cold-rolled steel sheet or steel strip not only has good elongation but also has good hole expanding performance, and is particularly suitable for automobile body. Manufacture of structural parts.
  • ultra-high-strength dual-phase steel has become the first choice for the automotive industry, because this advanced high-strength steel can effectively reduce the weight of the car body and improve safety.
  • High-strength steel sheets require not only good elongation in the automobile manufacturing process, but also high requirements for local forming ability, that is, high requirements for hole expansion ratio and bending performance.
  • the traditional cold-rolled dual-phase steel has a low yield ratio and has a certain drawing forming ability.
  • due to insufficient local formability it is prone to localized when manufacturing high-strength steel parts including deformation modes such as bending and reaming.
  • the band structure of steel is mainly caused by component segregation, and the segregation occurs in the solidification process of molten steel. Firstly, the content of solidified molten steel and the content of subsequent precipitated components are different. The concentration of alloying elements in molten steel will become higher and higher, eventually resulting in The content of the first solidified portion and the post-solidified portion of the solidified structure is very different. The region where the composition is segregated is deformed and elongated during the hot rolling to finally form a band structure.
  • the band structure usually contains high alloying elements, and it is difficult to eliminate due to the difficulty in diffusion of these alloying elements. The enrichment of the alloying elements attracts carbon in the same region, resulting in hardening of the duplex steel after banding.
  • the brittle martensite is more harmful to local deformation performance, and has better hole-expanding performance and cold-bending performance, and is prone to cracking during forming. Improving the uniformity of the structure and improving the local formability of the high-strength dual-phase steel is the key to obtaining a balanced dual-phase steel.
  • US Patent No. 20050167007A1 describes a method for producing a high-strength steel sheet, which is chemically formed. Divided into: 0.05-0.13% C 0.5-2.5% Si 0.5-3.5% Mn 0.05-1% Cr, 0.05-0.6% Mo, ⁇ 0.1% Al, ⁇ 0.005% S, ⁇ 0.01% N, ⁇ 0.03% P, 0.005-0.05% Ti or 0.005-0.05% Nb or 0.005-0.2% V is added. The steel is hot rolled at an Ar3 temperature or higher, taken up at 450-700 ° C, annealed, cooled and quenched from 700-600 ° C at a cooling rate of 100 ° C / s, and then tempered between 180-450 ° C. Finally, a high-strength steel having a tensile strength of 780 MPa and a hole expansion ratio higher than 50% was obtained.
  • Japanese Patent Laid-Open No. 11-350038 describes a steel of 980 MPa with good ductility and formability, and its composition is designed to be C: 0.1-0.15%, Si: 0.8-1.5%, Mn: 1.5-2.0%, P: 0.01- 0.05%, S ⁇ 0.005%, Sol Al: 0.01-0.07%, N: ⁇ 0.01%, Nb: 0.001-0.02%, V: 0.001-0.02%, Ti: 0.001-0.02% one or more.
  • the object of the present invention is to provide a high-formability cold-rolled ultra-high-strength steel plate, a steel strip and a method for producing the same, the tensile strength of the cold-rolled ultra-high-strength steel plate and the steel strip is ⁇ 980 MPa, and the strong plastic product is tensile strength ⁇
  • the elongation is ⁇ 17000
  • the hole expansion ratio is ⁇ 45%
  • the performance is balanced
  • the thickness ranges from 0.8 to 2.3 mm.
  • the steel is characterized by uniform distribution of the structure and small difference in hardness between the phases.
  • the main structures in the steel are ferrite, bainite, martensite and retained austenite.
  • the bainite in the steel is free of carbides, or only fine carbides are precipitated inside the bainite, and no interfacial carbides are precipitated. Compared with other cold-rolled steel sheets of the same grade, the performance is extended. Higher elongation or better hole expansion ratio or lower yield ratio, that is, more balanced mechanical properties, especially suitable for the formation of various types of automotive safety parts.
  • the steel of the invention has the advantages of good drawing performance and good hole expanding performance for high-strength steel forming, and can obtain higher elongation, lower yield ratio and expansion through proper composition design and process design.
  • the excellent comprehensive mechanical properties of the porosity are significantly superior to the existing steel grades in at least one characteristic, so that they have the advantage of balanced performance.
  • the invention adopts a higher carbon content design + high Si content design than the usual 980 MPa grade high strength steel, and a Mn content design which is equivalent to or slightly higher than the usual 980 MPa grade high strength steel.
  • the design of C, Si and Mn constitutes the basis of the composition design of the present invention: since the C content is significantly higher than that of the general 980 MPa grade high strength steel, it is easy to obtain higher retained austenite under the joint action of Si, Mn and the process. Thereby achieving a higher elongation.
  • the design of high silicon combined with reasonable technology, not only facilitates the acquisition of more retained austenite, but also facilitates the diffusion of C from bainite into austenite, thereby reducing the carbon content of bainite.
  • Alloying elements such as Mo, B, Ti, and Nb and microalloying elements are also added to the steel of the present invention.
  • Mo, B, Ti, and Nb and microalloying elements are also added to the steel of the present invention.
  • the addition of molybdenum element enhances the strength of the steel on the one hand, and promotes the fine precipitation of molybdenum and titanium by the design process in the hot rolling process, preferably in the interphase precipitation. These precipitates can be improved in the ferrite grains.
  • the hardness of the ferrite reduces the difference in hardness between the soft and hard phases, but does not substantially lower the elongation.
  • a small amount of zirconium is added to refine the grains of the prior austenite, and to reduce the concentration of the impurity elements at the grain boundaries. Adding B improves the tendency of P to segregate at grain boundaries. Further improvement of the plasticity and toughness of ultra high strength steel.
  • Ti and Nb can not only play the role of conventional grain refinement, but also cooperate with Mo to form phase-to-phase dispersion precipitation, which is more conducive to tissue uniformity, higher hole expansion ratio, and less reduction in elongation.
  • the high-formity cold-rolled ultra-high-strength steel sheet and steel strip of the present invention have a composition weight percentage of C: 0.15-0.35%, Si: 1.0-2.0%, Mn: 1.6-2.6%, and Mo: 0.1. -0.4%, P ⁇ 0.02%, S ⁇ 0.004%, N ⁇ 0.005%, Nb: 0.015-0.04%, Ti: 0.02-0.06%, Al: 0.015-0.045%, B: 0.0003-0.001%, and, B ⁇ P% / 30, the rest are Fe and inevitable impurities.
  • the weight percentage of the steel of the present invention is: C: 0.17-0.32%, Si: 1.2-1.8%, Mn: 1.8-2.5%, Mo: 0.15-0.4%, P: ⁇ 0.012%, S: ⁇ 0.002 %, N: ⁇ 0.005%, Nb: 0.015-0.04%, Ti: 0.02-0.06%, Al: 0.015-0.045%, B: 0.0003-0.001%, and B ⁇ P% / 30, and the remaining Fe and inevitable impurities .
  • the steel component of the present invention may further comprise Zr: 0.005 to 0.015% by weight.
  • the cold-rolled ultra-high-strength steel plate and the steel strip of the invention have tensile strength ⁇ 980 MPa, strong plastic product, that is, tensile strength ⁇ elongation ⁇ 17,000, and the hole expansion ratio is ⁇ 45%.
  • the cold rolled ultra-high strength steel plate and the steel strip of the invention have the characteristics of ferrite grain diameter ⁇ 10 ⁇ m, and the main structure is ferrite, bainite, martensite and residual austenite of less than 10% by volume. .
  • the carbon content is selected to be between 0.15 and 0.35%. If it is less than 0.15%, the strength is affected, and the amount and stability of austenite formation are lowered. If it is higher than 0.35%, the martensite hardness is too high. It is not conducive to the hole expansion rate, and the carbon equivalent is too high, which affects the welding performance, thus limiting the application.
  • Si plays a role in increasing the elongation in steel. Si has a great influence on the microstructure of steel, promoting the purification of ferrite and the formation of retained austenite. If less than 0.8%, the amount and amount of retained austenite formed are low, affecting the elongation of steel; if it is higher than 2.0%, it will bring other metallurgical quality defects, under the design premise of the present invention, Very necessary.
  • Mn Improves the hardenability of steel and effectively increases the strength of steel.
  • the content of Mn is 1.6-2.6%, the strength of steel below 1.6% is not enough, and the mechanism of promoting the formation of retained austenite is difficult to function; above 2.6%, the strength is too high, and segregation is easy to occur.
  • Mo can improve the hardenability of steel, effectively improve the strength of steel; Mo improves the distribution of carbides, and with appropriate hot rolling process, can form phase precipitation with Ti, improve the hardness of ferrite and improve the uniformity of the structure. It is good to increase the reaming rate. Adding 0.1-0.4% of Mo, less than 0.1% of Mo, the effect is not obvious, the carbide precipitation density is insufficient, higher than 0.4%, resulting in excessive yield strength.
  • Ti 0.02-0.04%, which acts to fix nitrogen and refine grains.
  • Ti and Mo combine to precipitate composite carbides. Especially when the hot rolling process is appropriate, it can obtain dispersed and fine phase precipitation. The hardness of the ferrite is increased, and it is not easy to roughen, and the hole expansion ratio can be improved better.
  • B It can improve the hardenability of steel and effectively increase the strength of steel; in the present invention, the addition amount of B is low, and is mainly used to reduce the tendency of intergranular segregation of P, so B: 0.0003-0.001%, and B ⁇ P%/30, B content and P content are related.
  • the P content is high, the B content is higher, which is beneficial to avoid the grain boundary segregation of P.
  • the B content is correspondingly reduced, because B is too high to have a greater influence on the strength.
  • Al It acts to deoxidize and refine grains in steel, and requires Al: 0.015-0.045%.
  • N It is an impurity element in steel and requires ⁇ 0.005%. Too high is likely to cause surface cracks or bubbles on the slab.
  • Nb It is a precipitation strengthening element, which plays a role in refining the grain and adjusting the strength. It is required to be distributed between 0.02 and 0.04%. If the concentration is too low, the strength increase is not obvious, and if it is too high, the plasticity decreases more. Nb refines grains and has certain benefits for tissue uniformity.
  • the invention seeks to reduce macrosegregation and microsegregation of S and P in steel in the manufacturing process.
  • the continuous casting process uses relatively rapid cooling, and the water spray per kilogram of steel is ⁇ 0.65 liters of water to refine the as-cast structure and reduce the degree of local segregation.
  • the water spray termination temperature is ⁇ 800 °C. This process facilitates obtaining a uniform as-cast structure.
  • the hot rolling process adopts 1100-1250 °C reheating and Ar3 or more finishing rolling, and adopts the first air cooling and water cooling cooling mode to ensure a certain slow cooling holding time between 780-800 ° C, thereby obtaining regular arrangement fine precipitation of phase precipitation. phase.
  • Annealing uses Ac3+30 °C above the holding temperature, using a higher primary cooling temperature and a higher rapid cooling start temperature to limit the formation of ferrite too high, or the redistribution of C in the high temperature zone is too full, Avoid excessive hardness due to excessive ferrite phase or soft ferrite phase.
  • Fast cooling requires cooling at a cooling rate of 40-120 ° C / s to between 200-400 ° C to ensure the necessary strength; at 200-400 ° C Tempering between, giving the opportunity to form retained austenite and bainite.
  • the final product has good elongation and hole expansion and thus good formability.
  • the method for producing a high-formability cold-rolled ultrahigh-strength steel sheet and a steel strip according to the present invention comprises the following steps:
  • the continuous casting billet adopts rapid cooling
  • the water spray per kg of steel is ⁇ 0.65 liters of water
  • the water spray termination temperature is ⁇ 800 °C.
  • v1 5-20 ° C / s cooling to the fast cold start temperature, fast cooling start temperature ⁇ 820-10 ⁇ v1, fast cooling at 40-120 ° C / s cooling to 200-450 ° C After tempering at 250-450 ° C for 100-400 s, it is then leveled by 0-0.3%.
  • the continuous casting billet adopts rapid cooling
  • the water spray amount per kilogram of steel is ⁇ 0.7 liters of water
  • the water spray termination temperature is ⁇ 800 °C.
  • step 2) hot rolling step the heating is carried out at 1100-1200 ° C, and the holding time is 0.8-1.2 hours, and hot rolling is performed at a temperature higher than Ar 3 .
  • air cooling is first performed, and a slow cooling state is maintained at 700-800 ° C for 10 s. Above, then rapid cooling, coiling temperature 500 ⁇ 600 ° C.
  • the speed of °C / s is cooled to 240-400 ° C, after tempering at 270-400 ° C for 100-400 s, and then 0-0.3% leveling.
  • the manufacturing process of the cold rolled high strength duplex steel strip of the invention is as follows:
  • the required alloy composition is obtained, and the content of S and P is minimized; the continuous casting billet adopts rapid cooling, and the segregation is minimized: the water spray amount per kg of steel is ⁇ 0.65 liters of water to refine the as-cast state. Organize and reduce the degree of local segregation, water spray termination temperature ⁇ 800 ° C.
  • 820-880 ° C heat preservation, higher soaking temperature, aim to obtain a more uniform structure, and cool to a rapid cooling start temperature at v1 5-20 ° C / s.
  • the quick cold start temperature is ⁇ 820-10 ⁇ v1.
  • Quick cold The starting temperature is related to the cooling rate of V1. If the cooling speed of V1 is faster, the formation of ferrite phase is less, and the diffusion of C is limited, the temperature of the rapid cooling can be lower. If V1 is low, ferrite is easy to form, and it is easy to soften, the rapid cold start temperature must be high.
  • the rapid cooling is cooled to 200-450 ° C at a speed of 40-120 ° C / s, after tempering at 250-450 ° C for 100-400 s, and then 0-0.3% leveling.
  • Fast cooling ensures sufficient strength and the tempering section ensures the formation of retained austenite and bainite. Leveling ensures the necessary shape.
  • the cold rolled ultra high strength steel sheet (belt) of the present invention has a thickness of 0.8 to 2.3 mm.
  • the performance characteristics of the cold-rolled ultra-high strength steel plate (belt) of the invention are: tensile strength ⁇ 980 MPa, high ductility (strong plastic product, tensile strength X elongation ⁇ 17000), high hole expansion ratio (reaming rate ⁇ 45%) With high drawability and high hole expansion ratio, it has balanced performance and is especially suitable for forming high-strength automotive parts.
  • the microstructure of the steel is characterized by fine and uniform microstructure, and the ferrite grain diameter is ⁇ 10 ⁇ m.
  • the main structure contained in the steel is ferrite, bainite, martensite and a small amount (less than 10% by volume). Austenite.
  • the ferrite grains are uniformly distributed in the steel, the bainite is precipitated in a short strip shape, and no carbides are precipitated between the bainite strips.
  • the retained austenite dispersion is present between the bainite strips or at the spaces between the ferrite grains. Martensite is dispersed throughout the tissue.
  • the steel of the invention has high strength and good formability, and has excellent elongation and hole expansion ratio, and has tensile strength ⁇ 980 MPa and high ductility (strong plastic product, ie tensile strength).
  • the forming of the strength of the automotive parts is well adapted to the needs of a variety of automotive parts manufacturing.
  • Table 1 shows the chemical composition of the steel embodiment of the present invention.
  • the manufacturing process of the steel embodiment of the present invention is shown in Table 2.
  • the strength of the steel of the present invention after smelting, hot rolling, cold rolling, annealing and flattening is as shown in Table 3. .
  • a high-strength cold-rolled steel sheet (belt) having a strength of 980 MPa or more can be produced, which has a good elongation and a good hole expansion ratio. It differs from the prior invention in terms of composition design, resource conservation, ease of manufacture, and ultimately obtained results and is superior to existing inventions.

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Abstract

一种高成形性的冷轧超高强度钢板、钢带及其制造方法，其成分重量百分比为：C 0.15-0.35%，Si 1.0-2.0%，Mn 1.6-2.6%，Mo 0.1-0.4%，P≤0.02%，S≤0.004%，N≤0.005%，Nb 0.015-0.04%，Ti 0.02-0.06%，Al 0.015-0.045%，B 0.0003-0.001%，且B≥P%/30，其余为Fe和不可避免的杂质。该钢板（带）抗拉强度≥980MPa，延伸率≥15%，扩孔率≥40%，性能均衡。

Description

一种高成形性的冷轧超高强度钢板、钢带及其制造方法 技术领域
本发明涉及一种高成形性的冷轧超高强度钢板、钢带及其制造方法，该冷轧钢板或钢带不仅具有较好的延伸率还有相当好的扩孔性能，特别适用汽车车身结构件的制造。
背景技术
汽车工业出于减重的需要，要求使用更高强度的钢板。其中，超高强度双相钢越来越成为汽车制造业的首选，因为这种先进高强钢，能有效减轻汽车车身重量，提高安全性。高强钢板在汽车制造过程中，不仅仅需要好的延伸率，同时对于局部成形能力要求很高，也即对扩孔率和弯曲性能要求较高。传统的冷轧双相钢，具有较低的屈强比，具备了一定的拉延成形能力，但由于局部成形性不足在制造包含弯曲和扩孔等变形方式的高强钢部件时，容易发生局部裂纹，从而影响整个零件的冲压效果，导致报废。文献研究表明，当双相钢扩孔率和弯曲性能偏低时，往往不能适应较为苛刻的成形条件，应用领域受到较大限制。高强度双相钢中一般含有较高的碳和合金元素，但较高的碳和合金元素容易导致铸造过程中发生成分偏析，造成后续的材料由于成分和组织的不均匀，造成局部变形能力下降，扩孔率和冷弯性差。钢中带状组织沿轧向分布，容易成为微观裂纹源，进一步降低钢的局部成形能力。
钢的带状组织主要是成分偏析引起的，偏析则发生于钢水凝固过程中，首先析出凝固的钢水成分和后续析出的成分含量不一样，钢水中的合金元素浓度会越来越高，最终造成凝固的组织中先凝固的部分和后凝固的部分合金元素含量差别非常大。成分偏析的区域在热轧过程中被变形拉长，最终形成带状组织。带状组织通常含有高的合金元素，并且由于这些合金元素扩散困难，很难消除，合金元素的富集吸引碳也富集在同样区域，造成双相钢淬火后形成呈带状分布的又硬又脆的马氏体，对局部变形性能危害较大，扩孔性能和冷弯性能均较低，成形过程中容易发生开裂。提高组织均匀性，提高高强双相钢的局部成形性是获得均衡型双相钢的关键。
美国专利US20050167007A1介绍了一种高强度钢板的制造方法，其化学成 分为：0.05-0.13％C 0.5-2.5％Si 0.5-3.5％Mn 0.05-1％Cr，0.05-0.6％Mo，≤0.1％Al，≤0.005％S，≤0.01％N，≤0.03％P，添加0.005-0.05％Ti或者0.005-0.05％Nb或者0.005-0.2％V。该钢经Ar3温度以上热轧，450-700℃卷取，退火后以100℃/s的冷速从700-600℃冷却淬火，然后在180-450℃之间回火。最终得到抗拉强度780Mpa的扩孔率高于50％的高强钢。
日本专利特开平11-350038介绍一种延性和成形性好的980MPa的钢，其成分设计为，C：0.1-0.15％，Si：0.8-1.5％，Mn：1.5-2.0％，P：0.01-0.05％，S≤0.005％，Sol Al：0.01-0.07％，N：≤0.01％，Nb：0.001-0.02％，V：0.001-0.02％，Ti：0.001-0.02％中的一种或以上。碳当量＝(C+Mn/6+Si/24)＝0.4-0.52，在Ar3以上热轧，500-650℃卷取，在Ac1-AC3之间保温，冷却到580-720℃，快冷到室温后，在230-300℃过时效。
中国专利号200810119823.0介绍的一种980MPa双相钢的制造方法，C：0.14-0.21％，Si：0.4-0.9％，Mn：1.5-2.1％，P：≤0.02％，S≤0.01％，Nb：0.001-0.05％，V：0.001-0.02％，经热轧冷轧后，在760-820℃间保温，冷速40-50℃/s，在240-320℃过时效180-300s。
以上专利多涉及980Mpa级别的高强钢，有的发明为普通的双相钢，特点是低屈服和适中的延伸率，具有拉延性能不错的优点；有发明为较高的扩孔率，但在获得高的扩孔率的同时，采用了较高的屈强比设计，拉延性能不足，扩孔率虽很高，但由于拉延性能不足，不适于有较高拉延要求的成形，其性能也不属于性能均衡的范畴。高强度钢许多的应用领域对拉延性能和扩孔性能均有较高的要求，如果仅仅拉延性能好，而扩孔性能不好，或者扩孔性能好，而拉延性能不好，则其应用领域就比较有局限性。
发明内容
本发明的目的是提供一种高成形性的冷轧超高强度钢板、钢带及其制造方法，该冷轧超高强钢板、钢带的抗拉强度≥980MPa，强塑积即抗拉强度×延伸率≥17000，扩孔率≥45％，性能均衡，其厚度范围0.8-2.3mm之间。该钢的特点是组织分布均匀，各相之间硬度差别小，钢中的主要组织为铁素体、贝氏体、马氏体和残余奥氏体。钢中的贝氏体无碳化物析出，或者只在贝氏体内部有细小的碳化物析出，无界面碳化物析出。性能方面和同级别的其它冷轧钢板相比，延 伸率更高或扩孔率更好或屈强比更低，即，具有更为均衡的力学性能，特别适用于各类汽车安全件的成形。
为达到上述目的，本发明的技术方案是：
本发明钢针对高强度钢成形既需要良好的拉延性能又需要良好的扩孔性能的特点，通过适当的成分设计和工艺设计，能得到较高的延伸率、较低的屈强比和扩孔率的优良综合力学性能，和现有同级别钢种相比，至少在一种特性方面明显占优，故具有性能均衡的优点。
本发明采用和通常980MPa级高强钢相比更高的含碳量设计+高Si含量设计，和通常980MPa级高强钢相当或略高的Mn含量设计。C、Si、Mn的设计构成了本发明成分设计的基础：由于C含量显著高于一般的980MPa级高强钢，在Si、Mn和工艺的共同作用下，容易获得更高的残余奥氏体，从而获得更高的延伸率。同时高硅的设计，配合合理的工艺，不仅有利于更多残余奥氏体的获得，还有利于C从贝氏体中扩散到奥氏体中去，从而降低贝氏体的的含碳量，从而有利于获得贝氏体内部无碳化物析出或者只有细小碳化物析出，边界没有碳化物析出。残余奥氏体的大量存在，有利于C从马氏体中向奥氏体中扩散，既提高了残余奥氏体稳定性，又降低了马氏体的含碳量，降低了马氏体的硬度，愈加对拉延性能和扩孔性能有利。采用适中的锰含量，以保证淬火性能，和碳元素一起，构成钢的强度的主要保证，同时如此设计范围的Mn和如此设计的Si、C，可以促进奥氏体的择优形成，进一步增加奥氏体的含量，有利于延性的提高。在如此高C、高Si和Mn含量的设计下，可以获得极低的屈强比、高的延伸率，但较低的扩孔率。
为了最终获得本发明钢低屈服、高延伸和高扩孔率的均衡性性能。本发明钢中还添加Mo、B、Ti、Nb等合金元素和微合金元素。添加钼元素，一方面提高钢的强度，另一方面在热轧工序利用设计的工艺促进钼和钛形成细微的析出，最好为相间析出，这些析出物在铁素体晶粒中，可以提高铁素体的硬度，减少软硬相之间的硬度差异，但基本不降低延伸率。添加微量的锆，细化原始奥氏体的晶粒，减轻杂质元素在晶界的浓度。添加B，改善P在晶界上偏聚的倾向。对于超高强度钢的塑性、韧性有进一步的改善效果。Ti、Nb不仅能起到常规的细化晶粒的效果，还能和Mo共同作用，形成相间的弥散析出，更加有利于组织均匀性，提高扩孔率，并且对于延伸率的降低较小。
具体的，本发明的高成形性的冷轧超高强度钢板、钢带，其成分重量百分比为：C：0.15-0.35％，Si：1.0-2.0％，Mn：1.6-2.6％，Mo：0.1-0.4％，P≤0.02％，S≤0.004％，N≤0.005％，Nb：0.015-0.04％，Ti：0.02-0.06％，Al：0.015-0.045％，B：0.0003-0.001％，且，B≥P％/30，其余为Fe和不可避免杂质。
优选的，本发明钢的成分重量百分比为：C：0.17-0.32％，Si：1.2-1.8％，Mn：1.8-2.5％，Mo：0.15-0.4％，P：≤0.012％，S：≤0.002％，N：≤0.005％，Nb：0.015-0.04％，Ti：0.02-0.06％，Al：0.015-0.045％，B：0.0003-0.001％，并且B≥P％/30，其余Fe和不可避免杂质。
进一步，本发明钢成分还可包含Zr：0.005-0.015％，重量百分比计。
本发明的冷轧超高强钢板、钢带的抗拉强度≥980MPa，强塑积即抗拉强度×延伸率≥17000，扩孔率≥45％。
本发明的冷轧超高强钢板、钢带的组织特点是：铁素体晶粒直径≤10微米，主要组织为铁素体、贝氏体、马氏体及体积百分数10％以下残余奥氏体。
在本发明钢的成分设计中：
C：提高钢的强度，提高马氏体的硬度，促进奥氏体的富碳，促进残余奥氏体的形成。因此，选择含碳量在0.15-0.35％之间，如果低于0.15％，强度受到影响，并且奥氏体的形成量和稳定性降低；如果高于0.35％，造成马氏体硬度过高，不利于扩孔率，同时碳当量过高，影响焊接性能，从而限制了应用。
Si：在钢中起到提高延伸率的作用。Si对钢的组织影响也很大，促进铁素体的纯净化和残余奥氏体的形成。如果低于0.8％，形成的残余奥氏体的量和量均较低，影响钢的延伸率；如果高于2.0％，会带来其它的冶金质量缺陷，在本发明的设计前提下，不是非常必要。
Mn：可提高钢的淬透性，有效提高钢的强度。选取Mn的含量为1.6-2.6％，低于1.6％钢的强度不够，并且促进残余奥氏体择优形成的机制难以发挥作用；高于2.6％，强度过高，也容易发生偏析。
C、Si、Mn的综合添加效果：通过高C+Si+Mn的设计，特别是较高的C含量水平下，可以促进较高的残余奥体体形成，获得低屈服、高延伸率的性能，详见前文所述。
Mo：可提高钢的淬透性，有效提高钢的强度；Mo改善碳化物的分布，配合适当的热轧工艺，可以和Ti共同形成相间析出，对提高铁素体的硬度，改善组织均匀性，提高扩孔率有好处。添加0.1-0.4％的Mo，低于0.1％的Mo，作用不明显，碳化物析出密度不足，高于0.4％，导致屈服强度过高。
Ti：0.02-0.04％，起到固定氮元素和细化晶粒的作用，Ti和Mo复合作用，析出复合碳化物，特别是在热轧工艺的适当时，可以获得弥散细小的相间析出，有效提高铁素体的硬度，并且不易粗化，可以更好低改善扩孔率。
B：可提高钢的淬透性，有效提高钢的强度；本发明中B的添加量较低，主要用于减轻P的晶间偏聚倾向，因此要求B：0.0003-0.001％，并且B≥P％/30，B含量和P含量进行关联，当P含量较高时，B含量较高，有利于避免P的晶界偏聚。当P含量低时，B含量相应降低，因为B过高会对强度影响较大。
Zr：0.0005-0.015％，细化原始奥氏体晶粒，减轻晶间杂质元素的浓度。
P：在钢中为杂质元素，要求≤0.02％。
S：在钢中为杂质元素，形成MnS严重影响扩孔率，要求≤0.004％。
Al：在钢中起到了脱氧作用和细化晶粒的作用，要求Al：0.015-0.045％。
N：在钢中为杂质元素，要求≤0.005％。过高容易导致板坯表面裂纹或气泡。
Nb：为析出强化元素，起到细化晶粒和调节强度的作用，要求分布在0.02-0.04％之间，过低对强度增加不明显，过高则塑性下降较多。Nb细化晶粒，对组织均匀性有一定益处。
本发明在制造工艺方面，力求降低钢中的S、P的宏观偏析和微观偏析。连铸工艺采用较快速的冷却，每公斤钢的喷水量≥0.65升水，以细化铸态组织和减轻局部偏析的程度，喷水终止温度≤800℃。这种工艺有利于获得均匀的铸态组织。热轧工艺采用1100-1250℃再加热、Ar3以上终轧后，采用先空冷后水冷的冷却方式，保证在780-800℃之间一定的缓冷保持时间，从而获得相间析出的规则排列微细析出相。退火采用Ac3+30℃以上的保温温度，采用较高的一次冷却温度和较高的快冷开始温度，以限制铁素体的形成量过高，或C在高温区的再分配过于充分，以避免铁素体相过多或铁素体相过软导致的硬度偏低。快冷要求以40-120℃/s的冷却速度，冷却到200-400℃之间，以保证必要的强度；在200-400℃ 之间回火，给残余奥氏体、贝氏体以形成的机会。最终产品拥有好的延伸率和扩孔率，从而具有好的成形性。
本发明高成形性的冷轧超高强度钢板、钢带的制造方法，其包括如下步骤：
1)冶炼、铸造
按上述成分冶炼、铸造，连铸坯采用快速冷却，每公斤钢的喷水量≥0.65升水，喷水终止温度≤800℃
2)热轧
采用1100-1250℃加热，0.6小时以上的保温时间，通过Ar3以上温度热轧，轧后首先空冷，在700-800℃之间保持缓冷状态5s以上，然后快速冷却，卷取温度500～600℃；
3)冷轧：40-65％压下率；
4)退火
820-880℃保温，以v1＝5-20℃/s冷却到快冷开始温度，快冷开始温度≥820-10×v1，快冷以40-120℃/s的速度冷却到200-450℃，经过250-450℃回火100-400s后，再经过0-0.3％平整。
优选的，连铸坯采用快速冷却，每公斤钢的喷水量≥0.7升水，喷水终止温度≤800℃。
优选的，步骤2)热轧工序中，采用1100-1200℃加热，0.8-1.2小时的保温时间，通过Ar3以上温度热轧，轧后首先空冷，在700-800℃之间保持缓冷状态10s以上，然后快速冷却，卷取温度500～600℃.
优选的，步骤4)退火工序中，保温温度830-860℃，以v1＝5-20℃/s冷却到快冷开始温度，快冷开始温度≥820-10×v1冷却速度，以40-120℃/s的速度冷却到240-400℃，经过270-400℃回火100-400s后，再经过0-0.3％平整。
本发明冷轧高强度双相带钢的制造工艺如下：
本发明冶炼和铸造过程中，获得要求的合金成分，尽量降低S、P的含量；连铸坯采用快速冷却，尽量减小偏析：每公斤钢的喷水量≥0.65升水，以细化铸态组织和减轻局部偏析的程度，喷水终止温度≤800℃。
退火工序中，820-880℃保温，较高的均热温度，目的获得更加均匀的组织，以v1＝5-20℃/s冷却到快冷开始温度。快冷开始温度≥820-10×v1。即快冷 开始温度的高低和V1的冷却速度有关，如果V1冷却速度较快，铁素体相的形成较少，C扩散有限，则快冷开始温度可以低一些。如果V1较低，铁素体易于形成，也易于软化，则快冷开始温度必须高。快冷以40-120℃/s的速度冷却到200-450℃，经过250-450℃回火100-400s后，再经过0-0.3％平整。快冷保证充分的强度，回火段保证残余奥氏体和贝氏体的形成。平整保证必要的板形。
本发明冷轧超高强钢板(带)厚度为0.8-2.3mm。
本发明冷轧超高强钢板(带)的性能特征为：抗拉强度≥980MPa，高延性(强塑积即抗拉强度X延伸率≥17000)，高扩孔率(扩孔率≥45％)，具备高拉延性和高扩孔率的特性，因此性能均衡，特别适合高强度汽车零件的成形。该钢的组织特点是组织细小、均匀，铁素体晶粒直径≤10微米，钢中含有的主要组织为铁素体、贝氏体、马氏体还含有少量(体积百分数10％以下)残余奥氏体。铁素体晶粒均匀分布于钢中，贝氏体呈短条状析出，贝氏体条之间无碳化物析出。残余奥氏体分散存在于贝氏体条之间或铁素体晶粒之间的空隙处。马氏体分散分布于组织中。
本发明的有益效果：
与现有技术相比，本发明钢的具有很高的强度和良好的成形性，其延伸率和扩孔率均很优良，其抗拉强度≥980MPa，高延性(强塑积即抗拉强度X延伸率≥17000)，高扩孔率(扩孔率≥45％)，具备高拉延性和高扩孔率的特性，实现了强度、延伸率和扩孔率的性能均衡性，特别适合高强度汽车零件的成形，很好地适应多种汽车零部件制造的需要。
本发明的最佳实施方式
下面结合实施例对本发明做进一步说明。
表1为本发明钢实施例的化学成分，本发明钢实施例的制造工艺如表2所示，本发明钢经冶炼、热轧、冷轧、退火和平整后得到的强度如表3所示。
从表3可以看出，按照本发明可以制造出强度980Mpa以上的高强度冷轧钢板(带)，其延伸率好，扩孔率好。在成分设计、资源节约和制造难易程度和最终获得的结果等各方面不同于现有发明且优于现有发明。
表1 本发明钢的化学成分(wt％)
Figure PCTCN2015070666-appb-000001
表2
Figure PCTCN2015070666-appb-000002
表3
Figure PCTCN2015070666-appb-000003
Figure PCTCN2015070666-appb-000004

Claims (13)

  1. 一种高成形性的冷轧超高强度钢板、钢带，其成分重量百分比为：C：0.15-0.35％，Si：1.0-2.0％，Mn：1.6-2.6％，Mo：0.1-0.4％，P≤0.02％，S≤0.004％，N≤0.005％，Nb：0.015-0.04％，Ti：0.02-0.06％，Al：0.015-0.045％，B：0.0003-0.001％，且，B≥P％/30，其余为Fe和不可避免杂质。
  2. 如权利要求1所述的高成形性的冷轧超高强度钢板、钢带，其特征是，其成分重量百分比为：C：0.17-0.32％，Si：1.2-1.8％，Mn：1.8-2.5％，Mo：0.15-0.4％，P：≤0.012％，S：≤0.002％，N：≤0.005％，Nb：0.015-0.04％，Ti：0.02-0.06％，Al：0.015-0.045％，B：0.0003-0.001％，并且B≥P％/30，其余Fe和不可避免杂质。
  3. 如权利要求1或2所述的高成形性的冷轧超高强度钢板、钢带，其特征是，还包括Zr：0.005-0.015％，重量百分比计。
  4. 如权利要求1或2所述的高成形性的冷轧超高强度钢板、钢带，其特征是，所述冷轧超高强钢板、钢带的抗拉强度≥980MPa，强塑积即抗拉强度×延伸率≥17000，扩孔率≥45％。
  5. 如权利要求1～4任一项所述的高成形性的冷轧超高强度钢板、钢带，其特征是，所述冷轧超高强钢板、钢带的组织特点是：铁素体晶粒直径≤10微米，主要组织为铁素体、贝氏体、马氏体及体积百分数10％以下残余奥氏体。
  6. 如权利要求1或2所述的高成形性的冷轧超高强度钢板带的制造方法，其特征是，包括如下步骤：
    1)冶炼、铸造
    按权利要求1或2的成分冶炼、铸造，连铸坯采用快速冷却，每公斤钢的喷水量≥0.65升水，喷水终止温度≤800℃；
    2)热轧
    采用1100-1250℃加热，0.6小时以上的保温时间，通过Ar3以上温度热轧，轧后首先空冷，在700-800℃之间保持缓冷状态5s以上，然后快速冷却，卷取温度500～600℃；
    3)冷轧：40-65％压下率；
    4)退火
    820-880℃保温，以v1＝5-20℃/s冷却到快冷开始温度，快冷开始温度≥820-10×v1，快冷以40-120℃/s的速度冷却到200-450℃，经过250-450℃回火100-400s后，再经过0-0.3％平整。
  7. 如权利要求6所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，连铸坯采用快速冷却，每公斤钢的喷水量≥0.7升水，喷水终止温度≤800℃。
  8. 如权利要求6所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，步骤2)热轧工序中，采用1100-1250℃加热，0.8-1.2小时的保温时间，通过Ar3以上温度热轧，轧后首先空冷，在700-800℃之间保持缓冷状态10s以上，然后快速冷却，卷取温度500～600℃.
  9. 如权利要求6所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，步骤4)退火工序中，保温温度830-860℃，以v1＝5-20℃/s冷却到快冷开始温度，快冷开始温度≥820-10×v1冷却速度，以40-120℃/s的速度冷却到240-400℃，经过270-400℃回火100-400s后，再经过0-0.3％平整。
  10. 如权利要求6所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，所述冷轧超高强度钢板、钢带的成分中还包括Zr：0.005-0.015％，重量百分比计。
  11. 如权利要求6或10所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，所述冷轧超高强度钢板、钢带的厚度为0.8-2.3mm。
  12. 如权利要求6或10或11所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，所述冷轧超高强钢板、钢带的抗拉强度≥980MPa，强塑积即抗拉强度×延伸率≥17000，扩孔率≥45％。
  13. 如权利要求6或10或11或12所述的高成形性的冷轧超高强度钢板、钢带的制造方法，其特征是，所述冷轧超高强钢板、钢带的组织特点是：铁素体晶粒直径≤10微米，主要组织为铁素体、贝氏体、马氏体及体积百分数10％以下残余奥氏体。
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Cited By (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN115181897A (zh) * 2021-04-02 2022-10-14 宝山钢铁股份有限公司 1280MPa级别低碳低合金超高强度双相钢及快速热处理制造方法
CN117925963A (zh) * 2023-12-14 2024-04-26 华北理工大学 一种超高强塑性超细贝氏体精轧螺纹钢及其制备方法
CN118207405A (zh) * 2024-05-21 2024-06-18 江苏永钢集团有限公司 10.9级高Cr免退火冷镦钢热轧盘条及其制造方法

Families Citing this family (19)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN106636899B (zh) * 2016-12-12 2018-08-03 东北大学 一种1000MPa级高扩孔型冷轧贝氏体钢的制造方法
CN108504956B (zh) * 2017-02-27 2020-07-28 宝山钢铁股份有限公司 高成型性冷轧超高强度复合钢板及其制造方法
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CN110129670B (zh) * 2019-04-25 2020-12-15 首钢集团有限公司 一种1300MPa级高强高塑性热冲压用钢及其制备方法
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JP7655221B2 (ja) * 2021-12-28 2025-04-02 Ｊｆｅスチール株式会社 鋼板およびその製造方法
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KR102845293B1 (ko) * 2022-12-29 2025-08-14 현대제철 주식회사 초고강도 냉연강판 및 그 제조방법
JP7655454B2 (ja) * 2023-03-28 2025-04-02 Ｊｆｅスチール株式会社 めっき鋼板、部材及びそれらの製造方法
EP4656759A1 (en) * 2023-03-28 2025-12-03 JFE Steel Corporation Steel sheet, member, and methods for producing same
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Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2006124798A (ja) * 2004-10-29 2006-05-18 Nippon Steel Corp 穴拡げ性に優れた高強度冷延薄鋼板およびその製造方法
US20080099109A1 (en) * 2006-10-31 2008-05-01 Hyundai Motor Company High-strength steel sheets with excellent formability and method for manufacturing the same
CN103857820A (zh) * 2011-10-07 2014-06-11 杰富意钢铁株式会社 成形性和耐冲击性优良的高强度热镀锌钢板及其制造方法
CN104040010A (zh) * 2012-01-13 2014-09-10 新日铁住金株式会社 冷轧钢板及冷轧钢板的制造方法

Family Cites Families (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP5418168B2 (ja) * 2008-11-28 2014-02-19 Ｊｆｅスチール株式会社 成形性に優れた高強度冷延鋼板、高強度溶融亜鉛めっき鋼板およびそれらの製造方法
CN102586688B (zh) * 2011-01-10 2016-03-30 宝山钢铁股份有限公司 一种双相钢板及其制造方法

Patent Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2006124798A (ja) * 2004-10-29 2006-05-18 Nippon Steel Corp 穴拡げ性に優れた高強度冷延薄鋼板およびその製造方法
US20080099109A1 (en) * 2006-10-31 2008-05-01 Hyundai Motor Company High-strength steel sheets with excellent formability and method for manufacturing the same
CN103857820A (zh) * 2011-10-07 2014-06-11 杰富意钢铁株式会社 成形性和耐冲击性优良的高强度热镀锌钢板及其制造方法
CN104040010A (zh) * 2012-01-13 2014-09-10 新日铁住金株式会社 冷轧钢板及冷轧钢板的制造方法

Cited By (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN115181897A (zh) * 2021-04-02 2022-10-14 宝山钢铁股份有限公司 1280MPa级别低碳低合金超高强度双相钢及快速热处理制造方法
CN115181897B (zh) * 2021-04-02 2023-07-11 宝山钢铁股份有限公司 1280MPa级别低碳低合金超高强度双相钢及快速热处理制造方法
CN117925963A (zh) * 2023-12-14 2024-04-26 华北理工大学 一种超高强塑性超细贝氏体精轧螺纹钢及其制备方法
CN118207405A (zh) * 2024-05-21 2024-06-18 江苏永钢集团有限公司 10.9级高Cr免退火冷镦钢热轧盘条及其制造方法

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