WO2024260403A1 - 一种超高强度钢带及其制造方法 - Google Patents

一种超高强度钢带及其制造方法 Download PDF

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WO2024260403A1
WO2024260403A1 PCT/CN2024/100303 CN2024100303W WO2024260403A1 WO 2024260403 A1 WO2024260403 A1 WO 2024260403A1 CN 2024100303 W CN2024100303 W CN 2024100303W WO 2024260403 A1 WO2024260403 A1 WO 2024260403A1
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Prior art keywords
steel strip
cooling
temperature
ultra
high strength
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English (en)
French (fr)
Inventor
张瀚龙
陈�光
郑涛
张玉龙
金鑫焱
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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 KR1020267000773A priority Critical patent/KR20260023021A/ko
Priority to EP24825291.8A priority patent/EP4733430A1/en
Priority to AU2024313824A priority patent/AU2024313824A1/en
Publication of WO2024260403A1 publication Critical patent/WO2024260403A1/zh
Priority to MX2025015570A priority patent/MX2025015570A/es
Anticipated expiration legal-status Critical
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • 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
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • 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
    • C21D1/26—Methods of annealing
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    • C21D11/00—Process control or regulation for heat treatments
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
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    • C21—METALLURGY OF IRON
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
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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
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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
    • 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
    • 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
    • C21D8/0247—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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—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
    • 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
    • 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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • 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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • 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
    • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
    • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
    • C23G1/02—Cleaning or pickling metallic material with solutions or molten salts with acid solutions
    • C23G1/08—Iron or steel
    • 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
    • C21D2211/004—Dispersions; Precipitations
    • 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

Definitions

  • the present invention relates to a high-strength steel and a manufacturing method thereof, and in particular to an ultra-high-strength steel with a coating and a manufacturing method thereof.
  • ultra-high-strength steel is increasingly widely used in automotive parts.
  • ultra-high-strength steel is currently produced in an integrated manner on a large scale, this places extremely high demands on the mechanical uniformity of the ultra-high-strength steel in the same roll.
  • the strength difference between the transverse and longitudinal directions of dual-phase steel with a tensile strength of 800MPa is about 8MPa, but when the tensile strength of dual-phase steel reaches 1000MPa, the strength difference between the transverse and longitudinal directions can be as high as about 40MPa.
  • One of the objects of the present invention is to provide an ultra-high strength steel strip having low anisotropy and high mechanical uniformity in the same coil while achieving a tensile strength of 1000 MPa.
  • the present invention provides an ultra-high strength steel strip, which contains Fe and inevitable impurity elements, and further contains the following chemical elements in the following mass percentages:
  • Its microstructure consists of uniformly dispersed granular bainite in the shape of a coral sea.
  • the present invention also provides an ultra-high strength steel strip, the mass percentage of each chemical element of which is:
  • the mass percentage of each chemical element satisfies at least one of the following items:
  • the C element In the ultra-high strength steel strip described in the present invention, the C element not only controls the microstructure phase change, but also forms alloy carbides with other alloy elements, thereby affecting the strength, formability and performance uniformity of the steel plate.
  • the C element content in the steel when the C element content in the steel is lower than 0.13%, on the one hand, the strength of the steel will not meet the target requirements, and on the other hand, it will also lead to insufficient formation of granular bainite; and if the C element content in the steel is higher than 0.20%, it is easy to generate martensitic structure, coarse cementite, and other forms of bainite (non-granular bainite, such as upper bainite), thereby deteriorating the performance and performance uniformity of the steel strip. Based on this, the present invention controls the mass percentage of C between 0.13 and 0.20%.
  • Si is an essential element for steelmaking and deoxidation, which has a certain solid solution strengthening effect and also has a certain influence on the formation of polygonal ferrite and bainite.
  • Si content in the steel when the Si content in the steel is lower than 0.15%, it is difficult to obtain a sufficient deoxidation effect; and when the Si content in the steel is higher than 0.5%, on the one hand, it is easy to form iron oxide scale or tiger stripe color difference, which is not conducive to the surface quality of the steel plate, and on the other hand, it will affect the formation of granular bainite and cause polygonal ferrite and carbon-free bainite to be easily formed in the steel, affecting the performance uniformity. Based on this, the present invention controls the mass percentage of Si between 0.15 and 0.5%.
  • Mn is one of the key controlling elements for the phase transformation of the steel structure.
  • the Mn element is too low, on the one hand, the strength of the steel will not meet the target requirements, and on the other hand, it will also lead to insufficient formation of granular bainite; and when the Mn content is too high, it will not only deteriorate the corrosion resistance and welding performance, promote the formation of non-granular bainite structures such as martensite, but also intensify the tendency of grain coarsening and the formation of banded structures or center segregation, reduce the formability of the steel, and degrade the uniformity and performance uniformity of the steel strip. Based on this, the present invention controls the mass percentage of Mn between 1.4 and 2.0%.
  • the B element is not only conducive to the formation of bainite in the steel, but also has a great influence on the strength and hardness of the steel plate. If the B element content in the steel is less than 0.001%, the strength of the steel will not meet the target requirements; and when the B element content in the steel is higher than 0.004%, brittle borides are easily generated, affecting the formability and uniformity of the steel plate. Based on this, the present invention controls the mass percentage of B to between 0.001 and 0.004%.
  • Al In the ultra-high strength steel strip described in the present invention, Al is added to the steel only as a deoxidizing element, which can remove the O element in the steel to ensure the performance and quality of the steel. Therefore, the present invention controls the mass percentage of Al to be between 0.01 and 0.04%.
  • Al is added to the steel in large quantities ( ⁇ 0.1%) as a ferrite forming element and an element that inhibits the precipitation of carbides, in the hope of bringing about solid solution strengthening, or by changing the phase transformation temperature (such as A1, A3), bainite formation kinetics and carbide precipitation kinetics through the addition of Al to change the phase transformation of the steel, forming residual austenite or carbon-free bainite, and ultimately improving the strength of the steel, it is not beneficial to improve the performance uniformity and isotropy of the steel strip, so there is no need to add a large amount of Al to avoid causing a substantial increase in cost and difficulty in steelmaking.
  • phase transformation temperature such as A1, A3
  • Cr and Mo In the ultra-high strength steel strip of the present invention, Cr and Mo can increase the hardenability of the steel strip, increase the incubation period of pearlite and ferrite, inhibit the formation of pearlite and ferrite, and make it easy to obtain bainite structure during cooling. Therefore, when the Cr and Mo content is too low, insufficient granular bainite formation will result. When the Cr and Mo content is too high, more martensite, tempered martensite and other structures will be easily formed, which will deteriorate the formability of the steel strip. On the other hand, it is easy to form banded structures in the steel strip. The weave morphology or center segregation deteriorates the uniformity of the steel strip and the uniformity of its performance. Therefore, in the present invention, Cr is controlled to 0.1-0.5% and Mo is controlled to 0.1-0.5%.
  • Ti, Nb and V are not intentionally added as common micro-alloying elements that are beneficial to the performance of steel, but are controlled as residual elements in steelmaking. Since these micro-alloying elements will form carbonitride precipitation from the formation of slabs during steelmaking continuous casting, especially the precipitation of larger carbonitrides such as TiN, and throughout the entire production process of the steel strip, it is extremely difficult to control the uniform, stable and on-demand precipitation of these carbonitrides, which will have an adverse effect on the uniformity of the performance of the steel strip. Therefore, in the present invention, Ti+Nb+V is limited to ⁇ 0.02%, and it is further preferably limited to Ti+Nb+V ⁇ 0.01%.
  • the mass percentage of each chemical element also satisfies: 2.7 ⁇ (Cr+Mo)/C ⁇ 3.3, where each chemical element in the formula is substituted with the value before the percentage sign of its mass percentage.
  • C, Cr, and Mo play a key role in the phase transformation of bainite
  • it is further preferably controlled to be 2.7 ⁇ (Cr+Mo)/C ⁇ 3.3. If (Cr+Mo)/C is too low, it is easy to cause the bainite phase area to be too small, the hardenability of the steel strip is insufficient, and during annealing production, too much ferrite, pearlite, upper bainite and other structures will be generated due to insufficient cooling rate. If (Cr+Mo)/C is too high, the hardenability of the steel strip will be too high and the martensite transformation point temperature will be high, resulting in excessive martensite in the steel strip.
  • the mass percentage of the inevitable impurity elements satisfies at least one of the following conditions: P ⁇ 0.012%, S ⁇ 0.004%, and N ⁇ 0.004%.
  • the impurity elements in the present invention are only P, S and N. If the technical conditions permit, it is hoped that the content is as low as possible. Further, the content can be controlled to be P ⁇ 0.012%, S ⁇ 0.004%, and N ⁇ 0.004%. In some embodiments, S ⁇ 0.003%.
  • the volume phase ratio of the granular bainite is ⁇ 95.0%.
  • the area of the granular bainite is ⁇ 5 ⁇ m 2 , and the aspect ratio is ⁇ 2:1.
  • the granular bainite is distributed in any area with a cross-sectional area ⁇ 50 2 ⁇ m in all areas 30 ⁇ m away from the upper and lower surfaces of the steel strip in the thickness direction.
  • granular bainite is distributed in any region with a cross-sectional area ⁇ 10 2 ⁇ m in all regions 20 ⁇ m away from the upper and lower surfaces of the steel strip in the thickness direction.
  • its microstructure also includes ferrite, and the volume phase ratio of the ferrite is 0.1-4.5%.
  • the microstructure of the ultra-high strength steel strip of the present invention may also contain retained austenite, other forms of bainite (such as acicular bainite), tempered martensite, titanium carbonitride (or niobium carbonitride, or vanadium carbonitride) and martensite.
  • bainite such as acicular bainite
  • tempered martensite titanium carbonitride (or niobium carbonitride, or vanadium carbonitride) and martensite.
  • the content of such structures is relatively small, the sum of the volume phase ratios does not exceed 0.5%, and the amount of the phase ratios satisfies the following: retained austenite>other forms of bainite>tempered martensite>titanium carbonitride (or niobium carbonitride, or vanadium carbonitride)>martensite.
  • the ultra-high strength steel strip described in the present invention has a tensile strength ⁇ 1000MPa and a yield strength ⁇ 780MPa; and the low anisotropy satisfies: the difference in yield strength between the transverse and longitudinal directions at the same position on the steel strip is ⁇ 25MPa, and the difference in tensile strength is ⁇ 20MPa; the mechanical uniformity of the same coil satisfies: for the same steel coil, along its length or width direction, the difference in yield strength at different positions in the same tensile direction is ⁇ 25MPa, and the difference in tensile strength is ⁇ 20MPa.
  • the ultra-high strength steel strip described in the present invention has a tensile strength ⁇ 1000MPa and a yield strength ⁇ 780MPa; and the low anisotropy satisfies: the difference in yield strength between the transverse and longitudinal directions at the same position on the steel strip is ⁇ 20MPa, and the difference in tensile strength is ⁇ 15MPa; the mechanical uniformity of the same coil satisfies: for the same steel coil, along its length or width direction, the difference in yield strength at different positions in the same tensile direction is ⁇ 20MPa, and the difference in tensile strength is ⁇ 15MPa.
  • the yield strength of the ultra-high strength steel strip of the present invention is ⁇ 800 MPa. In some embodiments, the yield strength of the ultra-high strength steel strip of the present invention is ⁇ 850 MPa. In some embodiments, the yield strength of the ultra-high strength steel strip of the present invention is ⁇ 880 MPa.
  • the ultra-high strength steel strip of the present invention has a tensile strength of 1000-1080 MPa and a yield strength of 780-950 MPa.
  • the ultra-high strength steel strip of the present invention has an elongation at break of ⁇ 10.0% and/or a hole expansion rate of ⁇ 50%.
  • the elongation at break of the ultra-high strength steel strip of the present invention is 10.0-15.0%, and/or the hole expansion ratio is 50-75%.
  • Another object of the present invention is to provide a method for manufacturing an ultra-high strength steel strip, by which a steel strip having a tensile strength of 1000 MPa and low anisotropy and high mechanical uniformity in the same roll can be produced.
  • the present invention also provides a method for manufacturing an ultra-high strength steel strip, which comprises the steps of:
  • Post-rolling cooling adopts fast cooling and slow cooling alternately for multi-stage cooling. Finally, for the area outside 100m of the head and tail of the steel strip, the main body of the steel strip is cooled to the main coiling temperature of 440°C ⁇ 520°C at a cooling rate of 5 ⁇ 100°C for coiling. For the area within 100m of the head and tail of the steel strip, the head and tail of the steel strip are cooled to the head and tail coiling temperature of 480°C ⁇ 560°C at a cooling rate of 5 ⁇ 25°C for coiling.
  • Annealing heat to 840-900°C at a heating rate of ⁇ 50°C/s and keep warm; then cool to 700-780°C at a cooling rate of 2-20°C/s and keep warm; then cool to 360-430°C at a cooling rate of 20-50°C and keep warm; then heat to 440-480°C at a heating rate of 5-30°C/s; then cool to 330-400°C at a cooling rate of 0.1-0.5°C/s; finally cool to room temperature at a cooling rate of 15-50°C/s and coil.
  • the present invention adopts the above-mentioned post-rolling cooling and coiling process in order to obtain a uniform cross-sectional shape of the steel strip, a good plate shape and a uniform precursor structure, so as to ensure that in the annealing stage the steel strip can obtain a coral sea-like uniformly dispersed granular bainite in the length, width and thickness directions, so as to ensure that the final steel strip can obtain low anisotropy and high coiling performance uniformity.
  • the main purpose of adopting a multi-stage cooling scheme is to ensure that the hot coil has uniform structure after coiling.
  • the input of cooling water will cause extremely efficient heat exchange and extremely rapid temperature changes, which can easily cause uneven structure performance in local areas of the hot coil.
  • due to the fluctuations in the cross-sectional shape and plate shape of the steel strip after fine rolling it is also easy to cause water accumulation in local areas of the hot coil, resulting in abnormal structure performance in local areas.
  • the main idea of the multi-stage cooling scheme of the present invention is to adopt a step-by-step fast and slow cooling alternating method to reduce the impact of cooling water input on the structure performance of the hot coil, gradually reduce the cooling rate, and must be followed by slow cooling after fast cooling to eliminate the influence of heat exchange fluctuations and water accumulation of cooling water.
  • the present invention adopts the method of coiling the main body of the steel strip at a low temperature and raising the coiling temperature within 100m from the head to the tail, mainly considering the influence of the environmental heat exchange after coiling on the microstructure and properties of the hot coil.
  • the effect of the coiling temperature is to form bainite in a relatively stable temperature range as much as possible: first, the coiling temperature of the main body is as low as possible, which greatly reduces the influence of environmental heat exchange on the dynamics of hot coil organization formation; second, the coiling temperature is slightly increased at the head and tail ⁇ 100m, which is also considered that the environmental heat exchange at the head and tail of the hot coil is high, which is easy to produce a large temperature drop.
  • 440°C ⁇ 520°C is the main temperature range for bainite formation. If the temperature is too low, it is easy to form martensite structure, and if the temperature is too high, it is easy to produce uneven heat exchange with the environment, thereby forming an uneven organization (pearlite, bainite or martensite may be formed in different areas).
  • the annealing process adopted in the present invention is mainly to ensure that the steel strip obtains uniformly dispersed granular bainite in the shape of a coral sea in the length, width and thickness directions.
  • the steel strip is heated to 840-900°C and kept warm at a heating rate of ⁇ 50°C/s, in order to ensure that the steel strip obtains the original austenite structure as uniform as possible. Then, it is cooled to 700-780°C and kept warm at a cooling rate of 2-20°C/s. On the one hand, it is to form a certain amount of ferrite structure at this time, thereby greatly reducing the possibility of martensite formation in the subsequent low-temperature bainite transformation process.
  • Too fast cooling rate is not conducive to the control of the uniformity of the steel strip organization, while too slow cooling rate is easy to form ferrite, pearlite and upper bainite.
  • the reheating stage and the slow cooling stage are designed to allow the steel strip to undergo phase transformation in the granular bainite transformation range over a relatively long period of time, thereby forming a coral sea-like uniformly dispersed granular bainite.
  • the heating temperature is controlled to be 1180-1280°C
  • the finishing rolling temperature is 870-970°C.
  • the heating temperature is controlled to be 1180-1280°C, on the one hand to ensure that the slab can be heated to a uniform temperature everywhere, and on the other hand, a higher heating temperature can also ensure that the subsequent finishing rolling temperature is 870-970°C. If the heating temperature is too low, the slab will be heated unevenly and the finishing rolling temperature will be too low. If the heating temperature is too high, the slab surface will be severely oxidized, and finally the slab surface structure will be abnormal after annealing.
  • the finishing temperature of the finishing rolling is controlled at 870-970°C. On the one hand, it is to avoid the transformation of austenite to ferrite during finishing rolling. On the other hand, a higher finishing temperature is beneficial to the cross-sectional size and Shape control: if the final rolling temperature is too low, ferrite transformation is likely to occur during or after finishing rolling, which is not conducive to the structure control during subsequent annealing; if the final rolling temperature is too high, more cooling water is needed to cool the steel strip during post-rolling cooling, and too much cooling water input is not conducive to the uniform control of the hot-rolled structure during coiling.
  • the difference between the center thickness and the thickness within 40 cm from the edge in the width direction of the steel strip section after finish rolling is controlled to be ⁇ 50 ⁇ m.
  • the difference between the center thickness and the thickness within 40 cm from the edge in the width direction of the steel strip section after fine rolling is controlled to be ⁇ 50 ⁇ m, mainly to cooperate with the subsequent cold rolling step to control the cross-sectional shape and plate shape of the steel strip during the final continuous annealing to be good, so as to reduce the uneven organizational properties of the finished coil caused by the fluctuation of the cross-sectional shape and plate shape of the steel strip during the final continuous annealing.
  • the temperature difference between the non-middle area and the middle area in the width direction of the steel strip is controlled to be ⁇ 30°C
  • the temperature fluctuation between the middle area and the non-middle area in the length direction is controlled to be ⁇ 15°C.
  • the process in this embodiment is to ensure the uniformity of the structural properties of the steel strip in the width and length directions after it is rolled by controlling the temperature fluctuation.
  • the cooling rate of rapid cooling is 50 to 200°C/s; the cooling rate of slow cooling is 5 to 30°C/s, preferably 5 to 25°C/s.
  • fast cooling is performed first and then slow cooling is performed, and multiple stages of cooling are performed in an alternating manner of fast cooling and slow cooling.
  • the stop cooling temperature of the first rapid cooling is 680-760° C. In some embodiments, the stop cooling temperature of the last slow cooling is 490-570° C. In some embodiments, except for the first rapid cooling, the temperature of the steel strip is reduced by 20-160° C., preferably 30-80° C., during each cooling.
  • the multi-stage cooling by alternating fast cooling and slow cooling specifically includes: in the first stage, the steel strip is cooled to 680-760°C at a cooling rate of 100-200°C/s; in the second stage, the steel strip is cooled to 600-670°C at a cooling rate of 5-30°C/s, preferably 5-25°C/s; in the third stage, the steel strip is cooled to 520-590°C at a cooling rate of 50-150°C/s; in the fourth stage, the steel strip is cooled to 490-570°C at a cooling rate of 5-30°C/s, preferably 5-25°C/s.
  • the selection of the step point temperatures of the above cooling sections is designed to be within the ferrite transformation range of the steel strip, so that the steel strip can preferentially transform a small amount of ferrite and pearlite at this stage to ensure the uniformity of the microstructure and properties of the hot coil, thereby greatly reducing the possibility of forming martensite during subsequent low-temperature coiling.
  • the cold rolling reduction ratio is controlled to be ⁇ 30%, such as 30-75%, and control the target thickness of the steel strip so that the difference between the thickness at the center point and the thickness at any position within 40 cm from the edge in the width direction of the steel strip section is ⁇ 30 ⁇ m.
  • the process parameters are set to control the cross-sectional shape and plate shape of the steel strip after cold rolling to be good, so as to reduce the uneven microstructure and performance of the finished coil caused by the fluctuation of the cross-sectional shape and plate shape of the steel strip during the final continuous annealing.
  • the heat preservation time is ⁇ 2min
  • the heat preservation time is >2min
  • the temperature difference between the non-middle area and the middle area in the width direction of the steel strip is controlled to be ⁇ 10°C at the same time, and the temperature fluctuation between the middle area and the non-middle area in the length direction of the steel strip is controlled to be ⁇ 5°C.
  • the process in this embodiment is also to ensure the uniformity of the structure and performance of the steel strip in the width and length directions after coiling by controlling the temperature fluctuation.
  • the material is heated to 840-900°C at a heating rate of ⁇ 50°C/s, such as 10-50°C/s, and kept warm for 1-4 minutes; then cooled to 700-780°C at a cooling rate of 2-20°C/s and kept warm for 10-40 seconds; then cooled to 360-430°C at a cooling rate of 20-50°C and kept warm for 2.5-10 seconds.
  • the present invention obtains ultra-high-strength steel with low anisotropy and high uniformity in the same roll through innovations in composition design, organizational regulation and manufacturing methods.
  • the tensile strength is ⁇ 1000MPa, and the yield strength is ⁇ 780MPa; and the low anisotropy satisfies: the difference in yield strength between the transverse and longitudinal directions at the same position on the steel strip is ⁇ 25MPa, and the difference in tensile strength is ⁇ 20MPa; the mechanical uniformity of the same coil satisfies: for the same steel coil, along its length or width direction, the difference in yield strength at different positions in the same tensile direction is ⁇ 25MPa, and the difference in tensile strength is ⁇ 20MPa.
  • the ultra-high strength steel strip of the present invention not only has the above-mentioned properties, but also has good formability, with a breaking elongation of ⁇ 10% and a hole expansion rate of ⁇ 50%.
  • the ultra-high strength steel strip of the present invention can be used for automotive parts with strict requirements on part size accuracy and service stability, such as automotive structural parts such as front seat slide rails in the automotive cabin system. It has a new design concept of "high, precise and advanced” refined, stable and differentiated products, and has good promotion prospects and application value.
  • FIG. 1 is a microstructure photograph of the ultra-high strength steel strip of Example 1 magnified 3000 times.
  • the ultra-high strength steel strips of Examples 1-7 and the comparative steel strips of Comparative Examples 1-2 of the present invention are prepared by the following steps:
  • Hot rolling control the heating temperature to 1180-1280°C and the finishing temperature to 870-970°C.
  • Post-rolling cooling adopts a multi-stage cooling method of alternating fast cooling and slow cooling.
  • the main body of the steel strip is cooled to the main coiling temperature of 440°C ⁇ 520°C at a cooling rate of 5 ⁇ 100°C for coiling.
  • the head and tail of the steel strip are cooled to the head and tail coiling temperature of 480°C ⁇ 560°C at a cooling rate of 5 ⁇ 30°C for coiling.
  • the above-mentioned multi-stage cooling by alternating fast cooling and slow cooling may specifically include: cooling the steel strip to 680-760°C at a cooling rate of 100-200°C/s in the first stage; cooling the steel strip to 600-670°C at a cooling rate of 5-30°C/s in the second stage; cooling the steel strip to 520-590°C at a cooling rate of 50-150°C/s in the third stage; cooling the steel strip to 490-570°C at a cooling rate of 5-30°C/s in the fourth stage, and making the steel strip temperature at the end of the fourth stage lower than the steel strip temperature at the end of the third stage.
  • the subsequent cooling of the steel strip to the coiling temperature is the fifth stage cooling, and making the coiling temperature at each location of the steel strip lower than the steel strip temperature at the corresponding position at the end of the fourth stage.
  • the target thickness of the steel strip is controlled so that in the width direction of the steel strip section, The difference between the thickness at the center point and the thickness at any position within 40 cm from the edge is ⁇ 30 ⁇ m.
  • the temperature difference between the non-middle area and the middle area in the width direction of the steel strip is controlled to be ⁇ 30°C
  • the temperature fluctuation between the middle area and the non-middle area in the length direction is controlled to be ⁇ 15°C.
  • 7 points of equal spacing temperature measurement and temperature control can be adopted in the width direction of the steel strip.
  • All temperatures in the hot rolling, cooling after rolling and coiling steps are the temperatures of the middlemost area of the 7-point equal spacing temperature measurement and temperature control, and the temperature difference between the other 6 areas and the middle area in the width direction can be controlled to be ⁇ 30°C.
  • All temperatures in the hot rolling, cooling after rolling and coiling steps refer to the average temperature of the middlemost area of the steel strip within the required length range, and the temperature fluctuation within the length range is ⁇ 15°C. When the length range is not specifically stated, it refers to the full length range of the steel strip.
  • Annealing heat to 840-900°C at a heating rate of ⁇ 50°C/s and keep warm for 1-4 min; then cool to 700-780°C at a cooling rate of 2-20°C/s and keep warm for 10-40 s; then cool to 360-430°C at a cooling rate of 20-50°C and keep warm for 2.5-10 s; then heat to 440-480°C at a heating rate of 5-30°C/s; then cool to 330-400°C at a cooling rate of 0.1-0.5°C/s; finally cool to room temperature at a cooling rate of 15-50°C/s to form a coil.
  • the holding time is ⁇ 2min
  • the heating temperature is ⁇ 870°C
  • the holding time is >2min
  • the temperature difference between the non-middle area and the middle area in the width direction of the steel strip is controlled to be ⁇ 10°C at the same time, and the temperature fluctuation between the middle area and the non-middle area in the length direction of the steel strip is controlled to be ⁇ 5°C.
  • 7 points of evenly spaced temperature measurement and temperature control can be used in the width direction of the steel strip.
  • All temperatures in the annealing step are the temperatures of the middlemost area of the 7-point evenly spaced temperature measurement and temperature control, and the temperature difference between the other 6 areas in the width direction and the middle area can be controlled to be ⁇ 10°C.
  • All temperatures in the annealing step refer to the average temperature of the middlemost area of the steel strip within the required length range, and the temperature fluctuation within the length range is ⁇ 5°C. When the length range is not specifically stated, it refers to the full length of the steel strip.
  • Table 1 lists the mass percentages of the chemical elements of the ultra-high strength steel strips of Examples 1-7 and the comparative steel strips of Comparative Examples 1-2.
  • Table 2-1, Table 2-2, Table 2-3 and Table 2-4 list the specific process parameters of the ultra-high strength steel strips of Examples 1-7 and the comparative steel strips of Comparative Examples 1-4 in the above process steps.
  • the microstructure detection method is as follows:
  • Preparation of metallographic specimens Use 4% nitric acid alcohol to pre-etch the polished metallographic specimens for about 10 seconds; further, the color metallographic specimen preparation method can be continued on this basis: the formula is: 1g sodium metabisulfite + 100ml water; and 4g picric acid + 100ml ethanol; the two solutions are mixed in equal proportions and allowed to stand, and then corroded for 30 to 40 seconds.
  • Observation Use an optical microscope or a scanning electron microscope to observe the microscopic structure, and take and analyze tissue photos.
  • the area of the granular bainite of the ultra-high strength steel strips of all embodiments is ⁇ 5 ⁇ m 2 , and the aspect ratio is ⁇ 2:1.
  • the granular bainite is distributed in any region with a cross-sectional area of ⁇ 50 2 ⁇ m in all regions 30 ⁇ m away from the upper and lower surfaces of the steel strips in the thickness direction.
  • the granular bainite is distributed in any region with a cross-sectional area of ⁇ 10 2 ⁇ m in all regions 20 ⁇ m away from the upper and lower surfaces of the steel strip in the thickness direction.
  • Figure 1 also shows the microstructure of the ultra-high strength steel strip of Example 1 magnified 3000 times.
  • the ultra-high strength steel strip has uniformly dispersed granular bainite in the shape of a coral sea.
  • the hole expansion rate of steel is determined by the hole expansion test.
  • the central hole of the specimen is expanded by pressing the specimen with a hole in the center into the die using a punch until necking or through cracks appear at the edge of the plate hole. Since the preparation method of the original hole in the center of the specimen and the corresponding quality of the edge of the original hole have a great influence on the hole expansion rate test results, the test and test method are carried out according to the hole expansion rate test method specified in the ISO/DIS 16630 standard, and the central original hole is in the form of a punched hole (corresponding to the processing method with the worst quality of the edge of the original hole).
  • the "strength difference” in Table 4-1 to Table 4-9 refers to the maximum and minimum differences of all test points in the longitudinal or transverse stretching direction of the full length and full width of the steel strip.
  • Comparative Example 2 meets the requirements of the present invention, the cooling rate in the manufacturing process is too high and does not meet the requirements of the present invention. Although the mechanical properties of the product meet the requirements, the uniformity of the same roll is poor and the strength difference between different positions is too large.

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Abstract

本发明公开了一种超高强度钢带，其含有Fe和不可避免的杂质元素，此外还含有质量百分含量如下的下述各化学元素：C：0.13～0.20％；Si：0.15～0.50％；Mn：1.4～2.0％；B：0.001～0.004％；Al：0.01～0.04％；Cr：0.1～0.4％；Mo：0.1～0.4％；Ti+V+Nb≤0.02％；其微观组织包括呈珊瑚海状的均匀弥散的粒状贝氏体。相应地，本发明还公开了该超高强度钢带的制造方法。本发明所述的超高强度钢带在抗拉强度达到1000MPa的同时，还具有低各向异性且高同卷力学均匀性。

Description

一种超高强度钢带及其制造方法 技术领域
本发明涉及一种高强钢及其制造方法，尤其涉及一种带镀层超高强钢及其制造方法。
背景技术
随着汽车工业的发展，车用零件的设计越发小巧、轻薄、精度高、且要求在双向甚至多个方向上的力学性能尽可能均匀，以确保零件高节拍生产、高精度的装配和高稳定性的服役。这就不仅要求材料具有超高的强度，良好的成形性能，还要有力学性能上较低的各向异性或较高的各向同性，更要有极高的性能均匀性。
现阶段，超高强钢越来越广泛的用于车用零件，同时由于超高强钢目前均采用规模化集成化生产，这就对超高强钢的同卷力学均匀性提出了极高的要求。
然而，在现有技术中，随着高强钢强度的提升达到1000MPa，要想获得低各向异性且高同卷力学均匀性的钢带难度极大。以公开的《塑性工程学报，2021，28(7)：124-130》中的数据为例，抗拉强度为800MPa级别的双相钢，其横纵向的强度差异约为8MPa，但当双相钢的抗拉强度达到1000MPa级别时，其横纵向的强度差异可高达约40MPa。
另外，公开号为CN104018092A，公开日为2017年3月8日，名称为“一种750MPa级别的高强度高力学性能均匀性的钢板及其制造方法”的中国专利文献，其公开的钢种的抗拉强度只有800MPa级别，横、纵、45°三向力学均匀性也只能达到<20MPa。
又例如，公开号为CN102397891A，公开日为2018年3月23日，名称为“一种提高钢带性能均匀性的方法”的中国专利文献，虽然其能够使得钢带宽度方向的强度均匀性较高，强度波动只有5～15MPa，但是这种钢种抗拉强度也只有约350MPa～600MPa。
由此可以看出，要想获得低各向异性(或高各向同性)且高同卷力学均匀性，同时强度达到1000MPa级别的超高强钢，是具有技术难度的。
发明内容
本发明的目的之一在于提供一种超高强度钢带，其在抗拉强度达到1000MPa的同时，具有低各向异性且高同卷力学均匀性。
为了实现上述目的，本发明提供了一种超高强度钢带，其含有Fe和不可避免的杂质元素，此外还含有质量百分含量如下的下述各化学元素：
C：0.13～0.20％；Si：0.15～0.50％；Mn：1.4～2.0％；B：0.001～0.004％；Al：0.01～0.04％；Cr：0.1～0.4％；Mo：0.1～0.4％；Ti+V+Nb≤0.02％；
其微观组织包括呈珊瑚海状的均匀弥散的粒状贝氏体。
相应地，本发明还提供了一种超高强度钢带，其各化学元素质量百分含量为：
C：0.13～0.20％；Si：0.15～0.50％；Mn：1.4～2.0％；B：0.001～0.004％；Al：0.01～0.04％；Cr：0.1～0.4％；Mo：0.1～0.4％；Ti+V+Nb≤0.02％；余量为Fe和不可避免的杂质元素。
进一步地，在本发明所述的超高强度钢带中，各化学元素质量百分含量满足下述各项的至少其中一项：
C：0.14～0.18％；
Ti+V+Nb≤0.01％。
在本发明所述的超高强度钢带中，各化学元素的设计原理如下所述：
C：在本发明所述的超高强钢带中，C元素不仅控制组织相变，也可以和其他合金元素形成合金碳化物，从而影响钢板的强度、成形性和性能均匀性。在本发明中，当钢中C元素含量低于0.13％，一方面会导致钢的强度达不到目标要求，另一方面也会导致粒状贝氏体形成不足；而若钢中C元素含量高于0.20％时，则容易生成马氏体组织、粗大的渗碳体、其他形态的贝氏体(非粒状贝氏体，如上贝氏体)，从而恶化钢带的性能和性能均匀性。基于此，本发明将C的质量百分比控制在0.13～0.20％之间。
Si：在本发明所述的超高强钢带中，Si元素是炼钢脱氧的必要元素，其具有一定的固溶强化作用，同时也会对多边形铁素体和贝氏体的形成有一定影 响。在本发明中，当钢中Si元素含量低于0.15％时，难以获得充分的脱氧效果；而当钢中Si元素含量高于0.5％时，一方面容易形成氧化铁皮或虎皮纹条状色差，不利于钢板的表面质量，另一方面会影响粒状贝氏体的形成，并导致钢中容易形成多边形铁素体和无碳贝氏体，影响性能均匀性。基于此，本发明将Si的质量百分比控制在0.15～0.5％之间。
Mn：在本发明所述的超高强钢带中，Mn是钢中组织相变关键的控制元素之一。当Mn元素过低时，一方面会导致钢的强度达不到目标要求，另一方面也会导致粒状贝氏体形成不足；而当Mn含量过高时，不仅会恶化耐腐蚀性能和焊接性能，促进马氏体等非粒状贝氏体组织的形成，同时也加剧晶粒粗化的趋势和带状组织或中心偏析形成的趋势，降低钢的成形性、劣化钢带的各相同性和性能均匀性。基于此，本发明将Mn的质量百分比控制在1.4～2.0％之间。
B：在本发明所述的超高强钢带中，B元素不仅有利于钢中贝氏体的形成，同时其还会对钢板的强度和硬度产生较大影响。若钢中B元素含量低于0.001％，则会导致钢的强度达不到目标要求；而当钢中B元素含量高于0.004％时，则容易生成脆性硼化物，影响钢板的成形性能和性能均匀性。基于此，本发明将B的质量百分比控制在0.001～0.004％之间。
Al：在本发明所述的超高强钢带中，Al元素仅作为脱氧元素加入钢中，其可以出去钢中的O元素，以保证钢的性能和质量。因此，本发明将Al的质量百分比控制在0.01～0.04％之间。虽然在一些现有技术中Al元素作为铁素体形成元素和抑制碳化物析出元素而大量(≥0.1％)加入钢中，以期望带来固溶强化，或通过Al的添加来改变相变温度(如A1，A3)、贝氏体形成动力学和碳化物析出动力学来改变钢材的相变，形成残余奥氏体或无碳贝氏体，并最终提高钢材强度，但对提高钢带的性能均匀性和各向同性并无益处，因此无需添加大量Al元素，以避免造成成本升高和炼钢制造难度的大幅增加。
Cr和Mo：在本发明所述的超高强钢带中，Cr和Mo可以增加钢带的淬透性，使珠光体和铁素体的孕育期增长，抑制珠光体的铁素体的形成，使冷却时易于得到贝氏体组织。因此Cr和Mo含量过低时，会导致粒状贝氏体形成不足，而当Cr和Mo含量过高时，则一方面容易形成较多的马氏体、回火马氏体等组织，恶化钢带的成形性能，另一方面容易导致钢带中形成带状组 织形态或中心偏析，从而恶化钢带的各相同性和性能均匀性。因此本发明中控制Cr：0.1～0.5％，Mo：0.1～0.5％。
Ti、Nb和V：在本发明所述的超高强钢带中，Ti、Nb和V并非是作为常见的对钢材性能有益的微合金元素进行有意添加的，而是作为炼钢的残余元素进行控制。因为这些微合金元素从炼钢连铸形成板坯开始，就会形成碳氮化物析出，尤其是TiN等尺寸较大的碳氮化物的析出，并贯穿于整个钢带的生产流程中，要想控制这些碳氮化物均匀、稳定、按需析出，难度极大，会对钢带的性能均匀性产生不利影响，因此在本发明中限制Ti+Nb+V≤0.02％，并进一步优选地限制Ti+Nb+V≤0.01％。
进一步地，在本发明所述的超高强度钢带中，其各化学元素质量百分含量还满足：2.7≤(Cr+Mo)/C≤3.3，式中的各化学元素代入其质量百分含量百分号以前的数值。
在本发明中，由于C、Cr、Mo对贝氏体的相变起到关键作用，可进一步优选地控制2.7≤(Cr+Mo)/C≤3.3。若(Cr+Mo)/C过低，则容易导致贝氏体相区过小，钢带淬透性不足，在退火生产时，会由于冷速的不足而生成过多的铁素体、珠光体、上贝氏体等组织，而若(Cr+Mo)/C过高，则会导致钢带淬透性过高且马氏体转变点温度偏高，导致钢带生成过多的马氏体。
进一步地，在本发明所述的超高强度钢带中，其中不可避免的杂质元素的质量百分含量满足下述各项的至少其中之一：P≤0.012％，S≤0.004％，N≤0.004％。
本发明中的杂质元素只要是P、S和N。在技术条件允许的情况下，希望其含量越低越好。进一步地，可以控制其含量P≤0.012％，S≤0.004％，N≤0.004％。在一些实施方案中，S≤0.003％。
进一步地，在本发明所述的超高强度钢带中，所述粒状贝氏体的体积相比例≥95.0％。
进一步地，在本发明所述的超高强度钢带中，所述粒状贝氏体的面积≤5μm2，长宽比≤2：1。
进一步地，在本发明所述的超高强度钢带中，在厚度方向上距钢带上下表面各30μm以外的全部区域内的任意一块截面面积≤502μm的区域内均分布有所述粒状贝氏体。
更进一步地，在本发明所述的超高强度钢带中，在厚度方向上距钢带上下表面各20μm以外的全部区域内的任意一块截面面积≤102μm的区域内均分布有粒状贝氏体。
这表示，在距钢带表面30μm以外的中间区域，基本上都分布有粒状贝氏体。
进一步地，在本发明所述的超高强度钢带中，其微观组织还包括铁素体，所述铁素体的体积相比例为0.1～4.5％。
除了粒状贝氏体和铁素体以外，本发明所述的超高强钢带的微观组织还可以含有残余奥氏体、其他形态的贝氏体(如针状贝氏体)、回火马氏体、碳氮化钛(或碳氮化铌、或碳氮化钒)和马氏体。但这种组织含量较小，体积相比例的总和不超过0.5％，且相比例的量满足残余奥氏体>其他形态的贝氏体>回火马氏体>碳氮化钛(或碳氮化铌、或碳氮化钒)>马氏体。
进一步地，本发明所述的超高强度钢带，其抗拉强度≥1000MPa，屈服强度≥780MPa；并且低各向异性满足：钢带上同一位置的横、纵向的屈服强度差≤25MPa、抗拉强度差≤20MPa；同卷力学均匀性满足：对于同一钢卷，沿其长度方向或宽度方向，相同拉伸方向上的不同位置的屈服强度差≤25MPa、抗拉强度差≤20MPa。
更进一步地，本发明所述的超高强度钢带，其抗拉强度≥1000MPa，屈服强度≥780MPa；并且低各向异性满足：钢带上同一位置的横、纵向的屈服强度差≤20MPa、抗拉强度差≤15MPa；同卷力学均匀性满足：对于同一钢卷，沿其长度方向或宽度方向，相同拉伸方向上的不同位置的屈服强度差≤20MPa、抗拉强度差≤15MPa。
在一些实施方案中，本发明所述的超高强度钢带的屈服强度≥800MPa。在一些实施方案中，本发明所述的超高强度钢带的屈服强度≥850MPa。在一些实施方案中，本发明所述的超高强度钢带的屈服强度≥880MPa。
在一些实施方案中，本发明所述的超高强度钢带的抗拉强度为1000～1080MPa，屈服强度为780～950MPa。
更进一步地，本发明所述的超高强度钢带，其断裂延伸率≥10.0％，并且/或者扩孔率≥50％。
在一些实施方案中，本发明所述的超高强度钢带的断裂延伸率为 10.0～15.0％，并且/或者扩孔率为50～75％。
本发明的另一目的在于提供一种超高强度钢带的制造方法，采用该方法可以制得抗拉强度达到1000MPa同时具有低各向异性且高同卷力学均匀性的钢带。
基于上述目的，本发明还提供了一种超高强度钢带的制造方法，其包括步骤：
冶炼和铸造；
热轧；
轧后冷却与卷取：轧后冷却采用快冷和慢冷交替的方式进行多段冷却，最后对钢带头尾100m以外的区域，以5～100℃的冷速将钢带主体冷却至主体卷取温度440℃～520℃进行卷取，对钢带头尾100m以内的区域，以5～25℃的冷速将钢带头尾冷却至头尾卷取温度480℃～560℃进行卷取；
酸洗和冷轧；
退火：以加热速率≤50℃/s的速率加热至840～900℃并保温；然后以2～20℃/s的冷速冷却至700～780℃并保温；然后以20～50℃的冷速冷却至360～430℃并保温；再以5～30℃/s的加热速率加热至440～480℃；然后以0.1～0.5℃/s的冷速冷却至330～400℃；最后以15～50℃/s的冷速冷却至室温成卷。
本发明采用上述轧后冷却与卷取工艺是为了获得均匀的钢带断面形状、良好的板型和均匀的前驱组织，以确保在退火阶段能够保证钢带在长度、宽度及厚度方向都获得珊瑚海状的均匀弥散的粒状贝氏体，以确保最终钢带能获得低各向异性和高成卷性能均匀性。
采用多段式冷却方案主要是保证卷取后的热卷组织均匀。在轧后冷却阶段，由于冷却水的投入会造成极高效的热量交换和极快速的温度变化，极易造成热卷局部区域组织性能不均匀，而且由于精轧后钢带断面形状和板型的波动，也容易造成热卷局部区域积水，造成局部区域组织性能异常，因此本发明多段式冷却方案的主体思路，就是采用阶梯式快慢冷却交替的方式，减少冷却水的投入对热卷组织性能的影响，冷速逐渐降低，且快冷后一定跟随慢冷以消除冷却水的热交换波动和积水的影响。
在最终的卷取温度上，本发明采用了钢带主体低温卷取，头尾100m以内提高卷取温度的方式，主要是考虑到卷取后环境热交换对热卷组织性能的影 响，并尽可能的在相对稳定的温度范围内形成贝氏体组织：首先，尽可能低的主体的卷取温度极大程度上减轻了环境热交换对热卷组织形成动力学的影响；其次，在头尾≤100m略微提高卷取温度，也是考虑到热卷头尾的环境热交换较高，容易产生较大温降。再次，440℃～520℃是贝氏体形成的主体温度范围，若温度过低，容易形成马氏体组织，而温度过高，又容易与环境产生不均匀的热交换，从而形成不均匀的组织(可能在不同的区域形成珠光体、贝氏体或马氏体)。
本发明采用的退火工艺，主要是为了保证钢带在长度、宽度及厚度方向都获得珊瑚海状的均匀弥散的粒状贝氏体。
基于此，首先以加热速率≤50℃/s的速率加热至840～900℃并保温，是为了保证钢带获得尽可能均匀的原始奥氏体组织。然后以2～20℃/s的冷速冷却至700～780℃并保温，一方面是为了此时形成一定量的铁素体组织，从而大幅减少后续低温贝氏体转变过程中马氏体形成的可能性，另一方面是为了降低快冷开始的温度，减少快冷区间的温降，以降低快冷时的热交换波动对组织性能的影响；然后通过快冷、加热和缓冷三个在贝氏体相区内的退火工艺设计，才可以形成珊瑚海状的均匀弥散的粒状贝氏体。其中，在快冷阶段，如果温度过低，则容易形成马氏体，如果温度过高，则容易形成上贝氏体，同时冷速的控制也很关键，过快的冷速不利于钢带组织均匀性的控制，而过慢的冷速则容易形成铁素体、珠光体和上贝氏体。再加热阶段和缓冷阶段的设计，是为了使钢带可以在一个较长的时间段内在粒状贝氏体转变区间内发生相变，从而形成珊瑚海状的均匀弥散的粒状贝氏体。
进一步地，在本发明所述的制造方法的热轧步骤中，控制加热温度为1180～1280℃，精轧终轧温度为870～970℃。
在该实施方式中，控制加热温度为1180～1280℃，一方面是为了保证板坯能够加热至各处温度均匀，另一方面较高的加热温度也可以保证后续的精轧终轧温度为870～970℃。若加热温度过低，会导致板坯加热不均匀，也会导致精轧的终轧温度过低，而若加热温度过高，则会导致板坯表面氧化严重，最终导致退火后板坯表面组织异常。
控制精轧终轧温度为870～970℃，一方面是为了避免在精轧时发生钢带的奥氏体向铁素体转变，另一方面，较高的精轧终轧温度有利于钢带断面尺寸及 形状控制，若终轧温度过低，则容易出现精轧中或精轧后的铁素体转变，不利于后续退火时的组织控制；若终轧温度过高，在轧后冷却时就需要更多的冷却水冷却钢带，而过多的冷却水投入不利于卷取时热轧组织的均匀控制。
进一步地，在本发明所述的制造方法的热轧步骤中，控制精轧后钢带断面的宽度方向上，中心点厚度与距边部40cm以内的厚度差≤50μm。
在该实施方式中，控制精轧后钢带断面的宽度方向上，中心点厚度与距边部40cm以内的厚度差≤50μm，主要是与后续的冷轧步骤配合，以控制最终连退时钢带的断面形状和板型良好，以减少最终连退时由于钢带的断面形状和板型的波动所造成的成品卷组织性能不均匀。
进一步地，在本发明所述的制造方法中，在热轧步骤、轧后冷却与卷取步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤30℃，长度方向上的中间区域与非中间区域的温度波动≤15℃。
该实施方式中的工艺是为了通过控制温度的波动，来保证钢带成卷后在宽度和长度方向上的组织性能均匀性。
在一些实施方案中，本文所述的轧后冷却步骤中，快冷的冷速为50～200℃/s；慢冷的冷速为5～30℃/s，优选5～25℃/s。
在一些实施方案中，本文所述的轧后冷却步骤中，先快冷后慢冷，以快冷和慢冷交替的方式进行多段冷却。
在一些实施方案中，首次快冷的停冷温度为680～760℃。在一些实施方案中，最后一次慢冷的停冷温度为490～570℃。在一些实施方案中，除首次快冷外，每一次冷却时钢带的温度降低20～160℃，优选30～80℃。
进一步地，在本发明所述的制造方法的轧后冷却步骤中，所述采用快冷和慢冷交替的方式进行多段冷却具体包括：第一段以100～200℃/s的冷速将钢带冷却至680～760℃；第二段以5～30℃/s、优选5～25℃/s的冷速将钢带冷却至600～670℃；第三段以50～150℃/s的冷速将钢带冷却至520～590℃；第四段以5～30℃/s、优选5～25℃/s的冷速将钢带冷却至490～570℃。
上述这些冷却段阶梯点温度的选择，是设计在钢带的铁素体转变区间，从而使钢带可以在此阶段优先转变出少量的铁素体、珠光体来确保热卷的组织性能均匀性，从而大幅减少了后续低温卷取时形成马氏体的可能性。
进一步地，在本发明所述的制造方法的冷轧步骤中，控制冷轧压下率≥ 30％，如30～75％，并控制钢带目标厚度，使得在钢带断面的宽度方向上，中心点厚度与距边部40cm以内的任意位置的厚度差≤30μm。
在该实施方式中，该工艺参数的设置是为了控制冷轧后的钢带的断面形状和板型良好，以减少最终连退时由于钢带的断面形状和板型的波动所造成的成品卷组织性能不均匀。
进一步地，在本发明所述的制造方法的退火步骤中，当进行加热和保温时，当加热温度≥870℃时，保温时间≤2min，而当加热温度<870℃时保温时间>2min。
进一步地，在本发明所述的制造方法中，在退火步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤10℃，控制钢带长度方向上的中间区域与非中间区域的温度波动≤5℃。
该实施方式中的工艺同样是为了通过控制温度的波动，来保证钢带成卷后在宽度和长度方向上的组织性能均匀性
进一步地，在本发明所述的制造方法的退火步骤中，以加热速率≤50℃/s、如10～50℃/s的速率加热至840～900℃并保温1～4min；然后以2～20℃/s的冷速冷却至700～780℃并保温10-40s；然后以20～50℃的冷速冷却至360～430℃保温2.5～10s。
本发明所述的超高强度钢带及其制造方法具有如下所述的优点以及有益效果：
本发明在保证化学元素成分和工艺相对简洁可控的前提下，通过成分设计、组织调控与制造方法的创新，获得了具有低各向异性和高同卷均匀性的超高强钢。
在一些实施方式中，其抗拉强度≥1000MPa，屈服强度≥780MPa；并且低各向异性满足：钢带上同一位置的横、纵向的屈服强度差≤25MPa、抗拉强度差≤20MPa；同卷力学均匀性满足：对于同一钢卷，沿其长度方向或宽度方向，相同拉伸方向上的不同位置的屈服强度差≤25MPa、抗拉强度差≤20MPa。
在一些实施方式中，本发明所述的超高强度钢带不仅具有上述性能，还具有良好的成形性能，其断裂延伸率≥10％，扩孔率≥50％。
本发明所述的超高强度钢带可用于对零件尺寸精度和服役稳定性等有严苛要求的汽车零部件，如汽车座舱系统中的前排座椅滑轨等汽车结构件，代表 了“高、精、尖”的精细化、稳定化、差异化产品的新设计理念，具有良好的推广前景和应用价值。
附图说明
图1为实施例1的超高强度钢带放大3000倍的微观组织照片。
具体实施方式
下面将结合具体的实施例对本发明所述的超高强度钢带及其制造方法做进一步的解释和说明，然而该解释和说明并不对本发明的技术方案构成不当限定。
实施例1-7和对比例1-2
本发明所述的实施例1-7的超高强度钢带和对比例1-2的对比钢带均采用以下步骤制得：
(1)按照表1所设计的化学成分配比进行冶炼和铸造。
(2)热轧：控制加热温度为1180～1280℃，精轧终轧温度为870～970℃。
在一些实施方式中，优选地控制精轧后钢带断面的宽度方向上，中心点厚度与距边部40cm以内的厚度差≤50μm。
(3)轧后冷却与卷取：轧后冷却采用快冷和慢冷交替的方式进行多段冷却，最后对钢带头尾100m以外的区域，以5～100℃的冷速将钢带主体冷却至主体卷取温度440℃～520℃进行卷取，对钢带头尾100m以内的区域，以5～30℃的冷速将钢带头尾冷却至头尾卷取温度480℃～560℃进行卷取。
在一些实施方式中，上述采用快冷和慢冷交替的方式进行多段冷却具体可以包括：第一段以100～200℃/s的冷速将钢带冷却至680～760℃；第二段以5～30℃/s的冷速将钢带冷却至600～670℃；第三段以50～150℃/s的冷速将钢带冷却至520～590℃；第四段以5～30℃/s的冷速将钢带冷却至490～570℃，并使得第四段末的钢带温度低于第三段末的钢带温度。后续将钢带进一步冷却到卷取温度则为第五段冷却，并使得钢带各处的卷取温度均低于第四段末对应位置的钢带温度。
(4)酸洗和冷轧：控制冷轧压下率≥30％。
在一些实施方式中，控制钢带目标厚度，使得在钢带断面的宽度方向上， 中心点厚度与距边部40cm以内的任意位置的厚度差≤30μm。
在一些优选地实施方式中，在热轧步骤、轧后冷却与卷取步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤30℃，长度方向上的中间区域与非中间区域的温度波动≤15℃。
在一些更具体的实施方式中，在热轧步骤、轧后冷却与卷取步骤的全过程中，可以在钢带宽度方向上采用7点等间距测温控温，热轧步骤、轧后冷却与卷取步骤中的所有温度为7点等间距测温控温的最中间区域的温度，可以控制宽度方向上其他6处区域与中间区域的温度差≤30℃。热轧步骤、轧后冷却与卷取步骤中的所有温度都是指钢带在所要求的长度范围内的最中间区域平均温度，在长度范围内的温度波动≤15℃，没有特定说明长度范围时，均指钢带全长范围内。
(5)退火：以加热速率≤50℃/s的速率加热至840～900℃并保温1～4min；然后以2～20℃/s的冷速冷却至700～780℃并保温10～40s；然后以20～50℃的冷速冷却至360～430℃保温2.5～10s；再以5～30℃/s的加热速率加热至440～480℃；然后以0.1～0.5℃/s的冷速冷却至330～400℃；最后以15～50℃/s的冷速冷却至室温成卷。
其中，当加热温度≥870℃时，保温时间≤2min，而当加热温度<870℃时保温时间>2min。
在一些实施方式中，在退火步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤10℃，控制钢带长度方向上的中间区域与非中间区域的温度波动≤5℃。
在一些更具体的实施方式中，在退火步骤的全过程中，可以在钢带宽度方向上采用7点等间距测温控温，退火步骤中的所有温度为7点等间距测温控温的最中间区域的温度，可以控制宽度方向上其他6处区域与中间区域的温度差≤10℃。退火步骤中的所有温度都是指钢带在所要求的长度范围内的最中间区域平均温度，在长度范围内的温度波动≤5℃，没有特定说明长度范围时，均指钢带全长范围内。
需要注意的是，实施例1-7的超高强度钢带所采用的化学成分设计以及相关制造工艺均满足本发明所设计的规范要求。相应地，对比例1-2的对比钢带所采用的化学成分设计以及相关制造工艺中均存在不满足本发明设计要求的 工艺参数。
表1列出了实施例1-7的超高强度钢带和对比例1-2的对比钢带的各化学元素的质量百分配比。
表1.(wt％，余量为Fe和除P、S、N以外其他不可避免的杂质)
表2-1、表2-2、表2-3和表2-4列出了实施例1-7的超高强度钢带和对比例1-4的对比钢带在上述工艺步骤中的具体工艺参数。
表2-1.
表2-2.
表2-3
表2-4

将所获得的实施例1-7的超高强度钢带分别进行取样，对其微观组织进行观察和检测，将微观组织观察结果列于下述表3中。
微观组织检测方法如下：
制备金相试样：采用4％硝酸酒精，对磨抛后的金相试样进行预侵蚀，侵蚀时间10s左右；进一步地，可在此基础上继续采用彩色金相试样制备方法：配方为：偏重亚硫酸钠1g+水100ml；与苦味酸4g+乙醇100ml；两种溶液等比例混合后静置，然后腐蚀30～40s。
观察：利用光学显微镜或扫描电子显微镜观察微观组织，并拍摄、分析组织照片。
表3.
此外，基于对各实施例的观察，所有实施例的超高强度钢带的粒状贝氏体的面积均≤5μm2，长宽比≤2：1。
实施例1-7的所有超高强度钢带，在厚度方向上距钢带上下表面各30μm以外的全部区域内的任意一块截面面积≤502μm的区域内均分布有所述粒状贝氏体。
实施例1-7的所有超高强度钢带，在厚度方向上距钢带上下表面各20μm以外的全部区域内的任意一块截面面积≤102μm的区域内均分布有所述粒状贝氏体。
此外，图1还显示了实施例1的超高强度钢带放大3000倍的微观组织。从图1可以看出，该超高强度钢带具有呈珊瑚海状的均匀弥散的粒状贝氏体。
此外，将所获得的实施例1-7的超高强度钢带和对比例1-2的对比钢带分别进行取样，对进行力学性能进行检测，以评测其各向异性和同卷均匀性，并在表4-1至4-9中分别列出了各实施例和对比例的各向异性和同卷均匀性。相关力学性能检测手段如下所述：
拉伸性能测试：采用GB/T228.1-2021《金属材料拉伸试验第1部分：室温试验方法》进行。
钢的扩孔率采用扩孔试验测定，通过使用凸模把中心带孔的试件压入凹模，使试件中心孔扩大，直到板孔边缘出现颈缩或贯穿裂纹为止。由于试件中心原始孔的制备方式及对应的原始孔边部质量对扩孔率测试结果存在较大影响，因此，试验及测试方法按ISO/DIS 16630标准中规定的扩孔率测试方法执行，中心原始孔采用冲压孔形式(对应为原始孔边部质量最差的加工方式)。
表4-1.实施例1
表4-2.实施例2

表4-3.实施例3
表4-4.实施例4

表4-5.实施例5
表4-6.实施例6

表4-7.实施例7
表4-8.对比例1

表4-9.对比例2
注：表4-1至表4-9中的“强度差”指的是钢带全长度方向以及钢带全宽度方向上纵向拉伸或者横向拉伸所有检测点的最大最小差值。
从上述表4-1至4-9可以看出，本发明所有实施例都满足抗拉强度≥1000MPa，屈服强度≥780MPa，断裂延伸率≥10％，扩孔率≥50％；同时各向 异性满足：钢带上同一位置的横、纵向的屈服强度差≤25MPa、抗拉强度差≤20MPa；同卷力学均匀性满足：对于同一钢卷，沿其长度方向或宽度方向，相同拉伸方向上的不同位置的屈服强度差≤25MPa、抗拉强度差≤20MPa。
而对比例1不仅C含量超出了设计要求的上限，而且不满足2.7≤(Cr+Mo)/C≤3.3，(Cr+Mo)/C仅为1.43，此时成分对冷速的要求极高，而在实际制造工艺中，在退火快冷段的冷速不满足要求，最终导致产品性能异常，不仅屈服强度和抗拉强度过低、扩孔率不足，同时各向异性较高，同卷均匀性较低。
对比例2虽然成分设计满足本发明要求，但是在制造工艺中，冷速过大不满足本发明要求，虽然产品力学性能满足要求，但同卷均匀性差，不同位置之间的强度差异过大。
需要说明的是，本案中各技术特征的组合方式并不限本案权利要求中所记载的组合方式或是具体实施例所记载的组合方式，本案记载的所有技术特征可以以任何方式进行自由组合或结合，除非相互之间产生矛盾。
还需要注意的是，以上所列举的实施例仅为本发明的具体实施例。显然本发明不局限于以上实施例，随之做出的类似变化或变形是本领域技术人员能从本发明公开的内容直接得出或者很容易便联想到的，均应属于本发明的保护范围。

Claims (15)

  1. 一种超高强度钢带，其含有Fe和不可避免的杂质元素，其特征在于，还含有质量百分含量如下的下述各化学元素：
    C：0.13～0.20％；Si：0.15～0.50％；Mn：1.4～2.0％；B：0.001～0.004％；Al：0.01～0.04％；Cr：0.1～0.4％；Mo：0.1～0.4％；Ti+V+Nb≤0.02％；
    其微观组织包括呈珊瑚海状的均匀弥散的粒状贝氏体。
  2. 如权利要求1所述的超高强度钢带，其特征在于，其各化学元素质量百分含量为：
    C：0.13～0.20％；Si：0.15～0.50％；Mn：1.4～2.0％；B：0.001～0.004％；Al：0.01～0.04％；Cr：0.1～0.4％；Mo：0.1～0.4％；Ti+V+Nb≤0.02％；余量为Fe和不可避免的杂质元素。
  3. 如权利要求1或2所述的超高强度钢带，其特征在于，其各化学元素质量百分含量满足下述各项的至少其中一项：
    C：0.14～0.18％：
    Ti+V+Nb≤0.01％。
  4. 如权利要求1或2所述的超高强度钢带，其特征在于，其Cr、Mo和C的质量百分含量还满足：2.7≤(Cr+Mo)/C≤3.3。
  5. 如权利要求1或2所述的超高强度钢带，其特征在于，其中不可避免的杂质元素的质量百分含量满足下述各项的至少其中之一：P≤0.012％，S≤0.004％，N≤0.004％。
  6. 如权利要求1或2所述的超高强度钢带，其特征在于，所述粒状贝氏体的体积相比例≥95％；优选地，所述粒状贝氏体的面积≤5μm2，长宽比≤2∶1；优选地，其微观组织还包括铁素体，所述铁素体的体积相比例为0.1～4.5％。
  7. 如权利要求1所述的超高强度钢带，其特征在于，在厚度方向上距钢带上下表面各30μm以外的全部区域内的任意一块截面面积≤502μm的区域内均分布有所述粒状贝氏体；优选地，在厚度方向上距钢带上下表面各20μm以外的全部区域内的任意一块截面面积≤102μm的区域内均分布有所述粒状贝氏体。
  8. 如权利要求1或2所述的超高强度钢带，其特征在于，其抗拉强度≥ 1000MPa，屈服强度≥780MPa、优选≥800MPa、更优选≥850MPa、进一步优选≥880MPa；并且低各向异性满足：钢带上同一位置的横、纵向的屈服强度差≤25MPa、抗拉强度差≤20MPa；同卷力学均匀性满足：对于同一钢卷，沿其长度方向或宽度方向，相同拉伸方向上的不同位置的屈服强度差≤25MPa、抗拉强度差≤20MPa；优选地，其断裂延伸率≥10％，并且/或者扩孔率≥50％。
  9. 如权利要求1-8中任意一项所述的超高强度钢带的制造方法，其特征在于，包括步骤：
    冶炼和铸造；
    热轧；
    轧后冷却与卷取：轧后冷却采用快冷和慢冷交替的方式进行多段冷却，最后对钢带头尾100m以外的区域，以5～100℃的冷速将钢带主体冷却至主体卷取温度440℃～520℃进行卷取，对钢带头尾100m以内的区域，以5～25℃的冷速将钢带头尾冷却至头尾卷取温度480℃～560℃进行卷取；
    酸洗和冷轧；
    退火：以加热速率≤50℃/s的速率加热至840～900℃并保温；然后以2～20℃/s的冷速冷却至700～780℃并保温；然后以20～50℃的冷速冷却至360～430℃并保温；再以5～30℃/s的加热速率加热至440～480℃；然后以0.1～0.5℃/s的冷速冷却至330～400℃；最后以15～50℃/s的冷速冷却至室温成卷。
  10. 如权利要求9所述的制造方法，其特征在于，所述热轧步骤，控制加热温度为1180～1280℃，精轧终轧温度为870～970℃；和/或，控制精轧后钢带断面的宽度方向上，中心点厚度与距边部40cm以内的厚度差≤50μm。
  11. 如权利要求9所述的制造方法，其特征在于，在热轧步骤、轧后冷却与卷取步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤30℃，长度方向上的中间区域与非中间区域的温度波动≤15℃；和/或，在退火步骤的全过程中，控制同一时刻下钢带宽度方向上非中间区域与中间区域的温度差≤10℃，控制钢带长度方向上的中间区域 与非中间区域的温度波动≤5℃。
  12. 如权利要求9所述的制造方法，其特征在于，在轧后冷却步骤中，快冷的冷速为50～200℃/s，慢冷的冷速为5～30℃/s、优选5～25℃/s；优选地，首次快冷的停冷温度为680～760℃，最后一次慢冷的停冷温度为490～570℃；优选地，除首次快冷外，每一次冷却时钢带的温度降低20～160℃，优选30～80℃；进一步优选地，所述采用快冷和慢冷交替的方式进行多段冷却具体包括：第一段以100～200℃/s的冷速将钢带冷却至680～760℃；第二段以5～30℃/s、优选5～25℃/s的冷速将钢带冷却至600～670℃；第三段以50～150℃/s的冷速将钢带冷却至520～590℃；第四段以5～30℃/s、优选5～25℃/s的冷速将钢带冷却至490～570℃。
  13. 如权利要求9所述的制造方法，其特征在于，在冷轧步骤中，控制冷轧压下率≥30％，并控制钢带目标厚度，使得在钢带断面的宽度方向上，中心点厚度与距边部40cm以内的任意位置的厚度差≤30μm。
  14. 如权利要求9所述的制造方法，其特征在于，在退火步骤中，当加热温度≥870℃时，保温时间≤2min，而当加热温度<870℃时，保温时间>2min。
  15. 如权利要求9所述的制造方法，其特征在于，在退火步骤中，以加热速率≤50℃/s的速率加热至840～900℃并保温1～4min；然后以2～20℃/s的冷速冷却至700～780℃并保温10-40s；然后以20～50℃的冷速冷却至360～430℃保温2.5～10s。
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Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20030217795A1 (en) * 2002-04-09 2003-11-27 Hitoshi Asahi High-strength steel sheet and high-strength steel pipe excellent in deformability and method for producing the same
CN108474057A (zh) * 2015-12-21 2018-08-31 安赛乐米塔尔公司 用于制造具有改善的延展性和可成形性的高强度钢板的方法和所获得的钢板
CN109930071A (zh) * 2019-04-09 2019-06-25 南阳汉冶特钢有限公司 一种建筑钢板及其生产方法
CN114086069A (zh) * 2021-03-04 2022-02-25 东北大学 一种含镁细晶粒热轧板带钢及其制备方法
KR20230059193A (ko) * 2021-10-25 2023-05-03 주식회사 포스코 강도와 저온 충격인성이 우수한 강재 및 그 제조방법

Family Cites Families (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN100392135C (zh) * 2005-06-30 2008-06-04 宝山钢铁股份有限公司 超高强带钢及其生产方法
CN114107785B (zh) * 2020-08-27 2022-10-21 宝山钢铁股份有限公司 一种具有超高屈强比的吉帕级贝氏体钢及其制造方法

Patent Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20030217795A1 (en) * 2002-04-09 2003-11-27 Hitoshi Asahi High-strength steel sheet and high-strength steel pipe excellent in deformability and method for producing the same
CN108474057A (zh) * 2015-12-21 2018-08-31 安赛乐米塔尔公司 用于制造具有改善的延展性和可成形性的高强度钢板的方法和所获得的钢板
CN109930071A (zh) * 2019-04-09 2019-06-25 南阳汉冶特钢有限公司 一种建筑钢板及其生产方法
CN114086069A (zh) * 2021-03-04 2022-02-25 东北大学 一种含镁细晶粒热轧板带钢及其制备方法
KR20230059193A (ko) * 2021-10-25 2023-05-03 주식회사 포스코 강도와 저온 충격인성이 우수한 강재 및 그 제조방법

Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN120796844A (zh) * 2025-09-08 2025-10-17 马鞍山钢铁有限公司 高强度座椅滑轨用钢及其生产方法

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