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WO2024061368A1 - 一种低点焊裂纹敏感性的带镀层超高强钢及其制造方法 - Google Patents

一种低点焊裂纹敏感性的带镀层超高强钢及其制造方法 Download PDF

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Publication number
WO2024061368A1
WO2024061368A1 PCT/CN2023/120965 CN2023120965W WO2024061368A1 WO 2024061368 A1 WO2024061368 A1 WO 2024061368A1 CN 2023120965 W CN2023120965 W CN 2023120965W WO 2024061368 A1 WO2024061368 A1 WO 2024061368A1
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Prior art keywords
strength steel
steel
coated ultra
manufacturing
strength
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Ceased
Application number
PCT/CN2023/120965
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English (en)
French (fr)
Inventor
潘华
雷鸣
金鑫焱
龚涛
蒋浩民
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Baoshan Iron and Steel Co Ltd
Original Assignee
Baoshan Iron and Steel Co Ltd
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Application filed by Baoshan Iron and Steel Co Ltd filed Critical Baoshan Iron and Steel Co Ltd
Priority to JP2025517397A priority Critical patent/JP2025531930A/ja
Priority to AU2023345395A priority patent/AU2023345395A1/en
Priority to KR1020257012041A priority patent/KR20250075594A/ko
Priority to EP23867653.0A priority patent/EP4592414A1/en
Publication of WO2024061368A1 publication Critical patent/WO2024061368A1/zh
Priority to MX2025003397A priority patent/MX2025003397A/es
Anticipated expiration legal-status Critical
Ceased legal-status Critical Current

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    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of 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
    • 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 by deformation combined with, or followed by, 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • 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
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/12—Aluminium or alloys based thereon
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/26—After-treatment
    • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
    • C23C2/36—Elongated material
    • C23C2/40—Plates; Strips
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C30/00—Coating with metallic material characterised only by the composition of the metallic material, i.e. not characterised by the coating process
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/06—Zinc or cadmium or alloys based thereon

Definitions

  • the invention relates to a high-strength steel and a manufacturing method thereof, in particular to a coated ultra-high-strength steel and a manufacturing method thereof.
  • the coating is mainly zinc-containing coating, such as hot-dip pure zinc, hot-dip galvanized iron and electroplated pure zinc.
  • Resistance spot welding has become the most important connection method in automotive welding production due to its advantages such as fast welding speed and low cost.
  • the joints often show good quasi-static mechanical properties, but a welding crack appears on the surface of the steel plate in contact with the electrode and its nearby area, and this welding crack is called spot welding.
  • Figure 1 schematically shows spot welding cracks in zinc-coated ultra-high-strength steel.
  • it is very difficult to suppress the formation of such cracks. Therefore, this crack has become an important obstacle and bottleneck for the application of zinc-coated ultra-high-strength steel in the automotive field.
  • the publication number is CN108015401A
  • the publication date is May 11, 2018
  • the Chinese patent document titled "Resistance spot welding method of galvanized high-strength steel with good joint performance” discloses an innovative spot welding process. This method ensures that the performance of the solder joint is not reduced while suppressing cracks on the spot welding surface.
  • the publication number is CN109385515A
  • the publication date is February 26, 2019,
  • the Chinese patent document titled "Multilayer Steel and Method for Reducing Liquid Metal Embrittlement” discloses a method to suppress spot welding cracks in high-strength steel. This method suppresses the occurrence of spot welding cracks by first decarburizing high-strength steel and then galvanizing it, so that the thickness of the decarburized layer is between 10 and 50 microns.
  • One of the purposes of the present invention is to provide a coated ultra-high-strength steel with low spot welding crack sensitivity.
  • the coated ultra-high-strength steel has very excellent quality and performance. While meeting the mechanical performance requirements of the head, it also has low spot welding crack sensitivity, which has good prospects for promotion and application.
  • the present invention provides a coated ultra-high-strength steel with low spot welding crack sensitivity, which includes a steel substrate and a zinc-containing coating on the surface of the steel substrate.
  • the steel substrate contains Fe and inevitable impurity elements,
  • the steel substrate also contains the following chemical elements in the following mass percentages:
  • the mass percentage of each chemical element of the steel substrate is:
  • each chemical element of the steel substrate further contains Mo: 0.10 to 2.00%.
  • the mass percentage content of C element is controlled at 0.10-0.30%.
  • the mass percentage of the C element can be further preferably controlled between 0.15% and 0.20%.
  • Si In the coated ultra-high-strength steel with low spot welding crack sensitivity according to the present invention, the solubility of silicon in carbides is extremely small, which can strongly inhibit the formation of cementite during the distribution process and promote the transfer of carbon to the residual It is enriched in austenite and improves the stability of retained austenite.
  • excessive silicon should not be added to steel. Too high a mass percentage of silicon will reduce the high-temperature plasticity of the steel, form stable oxides on the surface of the steel plate, and reduce the wettability of the steel plate.
  • silicon It is an element that significantly increases the tendency of cracks on the surface of solder joints. Based on this, in the coated ultra-high-strength steel with low spot welding crack sensitivity according to the present invention, the mass percentage content of Si element is specifically controlled at 0.30 to 0.90%.
  • the mass percentage of Si element can be further preferably controlled between 0.40% and 0.80%.
  • Mn In the coated ultra-high-strength steel with low spot welding crack sensitivity described in the present invention, manganese can expand the austenite phase area, reduce the Ac 3 , M s and M f points, and improve the austenite stability and steel The hardenability of the steel reduces the critical transformation rate, which is beneficial to preserving the retained austenite to room temperature. At the same time, manganese can also play a solid solution strengthening effect in the steel. When the Mn element content in the steel is too low, it will cause segregation. Ferrite and pearlite band structures are produced at low cooling rates. In addition, the mass percentage of manganese in steel should not be too high.
  • the mass percentage of manganese in steel When the mass percentage of manganese in steel is too high, it will intensify the grain coarsening trend, reduce the plasticity and toughness of the steel, and worsen the corrosion resistance. In particular, it will increase the enrichment of manganese in the superficial layer of the base metal under the coating, increase the surface crack sensitivity of the joint's solder joints, and worsen the welding performance. Based on this, in order to exert the beneficial effects of the Mn element, the mass percentage of the Mn element in the coated ultra-high-strength steel with low spot weld crack sensitivity according to the present invention is controlled at 1.00 to 2.20%.
  • the mass percentage of the Mn element can be further preferably controlled between 1.00% and 2.00%.
  • boron can significantly improve the hardenability of the steel.
  • boron easily segregates at grain boundaries, fills grain boundary defects, and reduces grain boundary energy, making it more difficult to nucleate new phases on austenite grain boundaries that are originally ferrite nucleation sites and increasing the stability of austenite. Strong, thus improving the hardenability; on the other hand, the addition of boron reduces the need for carbon, manganese and other elements that increase the hardenability of the steel.
  • boron element will reduce the segregation of manganese element to the grain boundary, thereby reducing the formation of manganese enrichment layer on the steel matrix under the coating, thereby significantly reducing the formation of spot welding cracks.
  • the more boron content the better.
  • a "boron phase” precipitation will be formed at the grain boundary, increasing the grain boundary energy.
  • the "boron phase” will serve as the core of the new phase, which will increase the nucleation speed, decrease the stability of austenite, and decrease the hardenability.
  • the inventors of this case have found that 0.0030% boron content is the inflection point that causes grain boundary embrittlement. Based on this, in order to give full play to the excellent effect of element B, the mass percentage of element B in the coated ultra-high strength steel with low spot welding crack sensitivity described in the present invention is controlled to be 0.001-0.003%.
  • the mass percentage of the B element can be further preferably controlled between 0.0015% and 0.003%.
  • the coated ultra-high-strength steel with low spot welding crack sensitivity when aluminum exists in a solid solution state, it can increase stacking fault energy, inhibit cementite precipitation and ⁇ to martensite transformation, and improve Austenite stability.
  • aluminum, carbon and nitrogen form fine dispersed refractory points, which can refine the grains.
  • the strengthening effect of aluminum is weaker than that of silicon, and its ability to stabilize austenite is also weaker than that of silicon.
  • the mass percentage of aluminum in steel is too high, it is easy to form a large number of oxide inclusions, which is not conducive to steelmaking and continuous casting. Therefore, in the coated ultra-high-strength steel with low spot weld crack sensitivity according to the present invention, the mass percentage content of the Al element is controlled between 0.30 and 1.00%.
  • molybdenum element can shift the C curve to the right and at the same time lower the Ms point, thereby improving the hardenability of the steel and improving the steel's hardenability. It not only improves the strength but also increases the elongation of steel; in addition, molybdenum is a strong carbide-forming element, which will form fine and dispersed MoC particles during smelting, and will be distributed in the form of a hard second phase during subsequent phase transformation. martensite, thus significantly improving the toughness of steel. More importantly, the inventor of this case found that the addition of molybdenum can significantly improve the resistance of steel to spot welding cracks.
  • the mass percentage of the Mo element is controlled between 0.10 and 2.00%. In some embodiments, the mass percentage of Mo element is controlled between 0.15% and 2.00%.
  • the mass percentage of Mo element can be further preferably controlled between 0.10% and 1.00%. In some embodiments, the mass percentage of the Mo element is controlled between 0.15 and 1.00% or between 0.20 and 1.00%.
  • the mass percentage content of each chemical element of the steel substrate satisfies at least one of the following items:
  • the mass percentage content of unavoidable impurity elements in the steel substrate satisfies: P ⁇ 0.01%, S ⁇ 0.005%.
  • P and S elements are both impurity elements in the ultra-high strength steel with coating described in the present invention.
  • the content of impurity elements in the ultra-high strength steel with coating should be reduced as much as possible.
  • P and S elements are both impurity elements.
  • P can play a solid solution strengthening role, inhibit the formation of carbides, and is beneficial to improving the stability of retained austenite, if the mass percentage of P is too high, it will Weakening grain boundaries, increasing material brittleness, and deteriorating welding performance, that is to say, the positive effect of P element is weaker than its negative effect. Therefore, it is preferable to control the mass percentage of P to P ⁇ 0.01%.
  • the S element in steel easily forms a low-melting eutectic at the grain boundary.
  • the mass percentage of the S element is controlled to S ⁇ 0.005%. .
  • the microstructure of the steel substrate is ferrite+martensite+retained austenite.
  • the volume proportion of ferrite is 25% to 45%; and/or the volume proportion of martensite is 45% to 65%.
  • the volume proportion of crystal grains with a size of 10 ⁇ m or less is ⁇ 85%, and the volume proportion of crystal grains with a size of 5 ⁇ m or less is ⁇ 55%.
  • the average grain size of the retained austenite is ⁇ 2 ⁇ m; and/or the average C content in the retained austenite is ⁇ 1.0%, such as 1.0-1.3%. .
  • the "low spot welding crack sensitivity" means that no cracks on the surface of the solder joint are produced when the spot welding process is used, or if cracks on the surface of the solder joint are produced, crack, the maximum length of the crack on the surface of the solder joint is less than 5% of the plate thickness, preferably less than 3% of the plate thickness, and more preferably less than 1% of the plate thickness.
  • its mechanical properties satisfy: its yield strength ⁇ 600MPa, such as ⁇ 700MPa, preferably ⁇ 750MPa; tensile strength ⁇ 980MPa, preferably ⁇ 1040MPa; uniform elongation ⁇ 13 %, preferably ⁇ 15.5%; elongation at break ⁇ 15%, preferably ⁇ 20%, more preferably ⁇ 22.5%.
  • its mechanical properties meet the following requirements: its yield strength is 600MPa ⁇ 850MPa, its tensile strength is 980MPa ⁇ 1150MPa, the uniform elongation is not less than 13%, and the fracture elongation is not less than 15%.
  • its mechanical properties meet: its yield strength is 755-845MPa, the tensile strength is 1040-1140MPa, the uniform elongation is 14.7-16.4%, and the fracture elongation is between Between 22.6 ⁇ 25.4%.
  • the coating is pure zinc coating, zinc-iron alloy coating, zinc-aluminum-magnesium coating or aluminum-zinc coating.
  • the coating weight is 30 to 120 g/m 2 , preferably 50 to 100 g/m 2 .
  • another object of the present invention is to provide a method for manufacturing coated ultra-high-strength steel.
  • the manufacturing method is simple to produce, and the obtained high-strength steel has excellent spot welding performance, especially anti-solder spot surface, under the same mechanical properties. The ability to crack is significantly improved.
  • the present invention proposes a manufacturing method of coated ultra-high-strength steel, which specifically includes the steps:
  • Hot rolling Control the thickness of the oxide scale on the surface of the steel strip after hot rolling to ⁇ 4 ⁇ m, and the mass percentage of FeO + Fe 3 O 4 in the oxide scale on the surface of the steel strip after hot rolling is ⁇ 50wt%;
  • Continuous annealing annealing at 800-920°C, then slowly cooling to 700-770°C at a cooling rate of 3-10°C/s; then cooling at 50-500°C/s (e.g., 100-500°C/s, 200-500°C/s, or 400-500°C/s) Cool quickly to 200-300°C; then heat to 360-460°C, keep warm for 50-600s; finally cool to room temperature;
  • 50-500°C/s e.g., 100-500°C/s, 200-500°C/s, or 400-500°C/s
  • the manufacturing method adopts a thin slab continuous casting process combined with a pickling or acid rolling process, and after continuous annealing and coating manufacturing, coated ultra-high-strength steel with low spot welding crack sensitivity can be obtained .
  • the coated ultra-high-strength steel produced by this manufacturing method can be welded using the conventional spot welding process of the car factory, and the maximum length of the crack on the surface of the welding point is less than 5% of the plate thickness. It has very low spot welding crack sensitivity and is very Excellent quality.
  • step (1) of the present invention since thin slab continuous casting is used, the rough rolling process can be omitted and the amount of hot rolling deformation can be reduced, thereby ensuring the steel plate performance in the subsequent steps (4) and (5). .
  • thin slab continuous casting since thin slab continuous casting is used in step (1), it can fully utilize the heat of the slab and reduce the energy consumption required for heating, thereby obtaining a more uniform ferrite or ferrite + pearlite structure, which in turn has It is beneficial to maintain a certain amount of fine-grained ferrite in the microstructure of the substrate of the finished product in step (6) and improve the uniformity of the structure.
  • step (2) the thickness of the oxide scale on the surface of the hot-rolled steel strip is controlled to ⁇ 4 ⁇ m, and the (FeO + Fe 3 O 4 ) in the oxide scale on the surface of the hot-rolled steel strip is ⁇ 50wt%, which is beneficial to the subsequent step ( 4), and has an important impact on the properties of the steel plate obtained after continuous annealing.
  • step (5) the inventor optimized the design of the continuous annealing process, which can form a homogenized austenite structure or austenite + ferrite by controlling the annealing temperature of 800 to 920°C for annealing.
  • the structure is then slowly cooled to 700-770°C at a cooling rate of 3-10°C/s to further adjust the ferrite content in the structure and obtain a certain proportion of ferrite, thereby improving the shaping of the material; and then using Cooling to 200 ⁇ 300°C at a speed of 50 ⁇ 500°C/s (that is, between M s (the starting temperature of martensite transformation) and M f (the end temperature of martensitic transformation)), at this time, the austenitic The body part is transformed into martensite, which can ensure that the steel obtains higher strength; then, it is heated to 360 ⁇ 460°C and kept for 50 ⁇ 600s.
  • High-strength steel has low spot welding crack sensitivity, that is, the maximum length of cracks on the surface of the welding spot is less than 5% of the plate thickness.
  • the zinc-containing coating when the zinc-containing coating is plated in step (6), the zinc-containing coating can be produced by, but is not limited to, hot-dip, electroplating, or vacuum plating techniques.
  • the coated ultra-high-strength steel designed by the present invention adopts the chemical composition design of carbon, silicon, manganese, boron, aluminum, and molybdenum and cooperates with the ferrite grain refinement. Therefore, during the continuous annealing process, while the nucleation points of austenite reverse transformation increase, the grain size can be further refined, so that the average grain size of the retained austenite stably maintained at room temperature is ⁇ 2 ⁇ m; the residual austenite can be stably maintained at room temperature.
  • the average C content in austenite is ⁇ 1.0%.
  • the thickness of the slab at the exit end of the thin strip continuous casting is controlled to be 50 to 58 mm.
  • step (1) the casting speed of the thin strip is controlled to be 2 to 5 m/min.
  • step (2) the slab is heated to 1200-1250°C.
  • step (3) the final rolling temperature is controlled to be 860-930°C, and the coiling temperature is controlled to be 450-600°C.
  • step (4) when pickling + cold rolling is used, the cold rolling deformation is controlled to be 40% to 60%.
  • step (5) the volume content of hydrogen in the reducing atmosphere in the continuous annealing furnace is controlled to 10-15%.
  • the step (5) The annealing process parameters meet at least one of the following:
  • Annealing temperature is 820 ⁇ 870°C
  • the coated ultra-high-strength steel with low spot weld crack sensitivity and its manufacturing method according to the present invention have the following advantages and beneficial effects:
  • the coated ultra-high-strength steel designed by the present invention its chemical element composition is based on C, Si, Mn, B, Al and Mo. By optimizing the ratio of each element, low spot welding crack susceptibility can be obtained. Durable coated ultra-high-strength steel.
  • the coated ultra-high-strength steel with low spot weld crack sensitivity described in the present invention has very excellent quality and performance. While meeting the user's requirements for the performance of the coated high-strength steel and the mechanical properties of the spot welded joint, it also has low points. Weld crack susceptibility.
  • the coated ultra-high-strength steel prepared using the technical solution of the present invention includes a steel substrate and a zinc-containing coating on the surface of the steel substrate. Its mechanical properties meet: its yield strength is 600MPa ⁇ 850MPa, its tensile strength is 980MPa ⁇ 1150MPa, and it extends uniformly. The rate is not less than 13%, and the elongation at break is not less than 15%; and, when spot welding is actually used for welding, if cracks on the surface of the solder joint occur, the maximum length of the crack on the surface of the solder joint is less than 5% of the plate thickness.
  • the manufacturing method designed by the present invention has a simple production process, and the obtained high-strength steel has significantly improved resistance to spot welding cracks under the same mechanical properties, and will have good application prospects in the production of safety structural parts for downstream users.
  • Figure 1 schematically shows spot welding cracks in zinc-coated ultra-high-strength steel.
  • the coated ultra-high-strength steels of Examples 1-28 and the comparative steels of Comparative Examples 1-4 according to the present invention are prepared by the following steps:
  • Hot rolling Control the oxide scale thickness on the surface of the steel strip after hot rolling to ⁇ 4 ⁇ m, and (FeO + Fe 3 O 4 ) in the oxide scale on the surface of the steel strip after hot rolling ⁇ 50wt%, and control the final rolling temperature to 860 ⁇ 930 °C, the coiling temperature is 450 ⁇ 600°C.
  • Continuous annealing Control the volume content of hydrogen in the reducing atmosphere in the continuous annealing furnace to 10 to 15%; anneal at 800 to 920°C, preferably the annealing temperature can be controlled to 820 to 870°C, and then 3 to 10 Slowly cool to 700-770°C at a cooling rate of °C/s to obtain a certain proportion of ferrite.
  • the final cooling temperature of slow cooling can preferably be controlled between 700-730°C; and then cool at a cooling rate of 50-500°C/s. Cool quickly to 200 ⁇ 300°C to partially transform austenite into martensite.
  • the final cooling temperature of rapid cooling can preferably be controlled between 250 ⁇ 300°C; then heat to 360 ⁇ 460°C after rapid cooling , preferably 400 to 430°C, holding for 50 to 600s, preferably holding for 180 to 300s; and finally cooled to room temperature.
  • the coating can be specifically selected as pure zinc coating, zinc-iron alloy coating, zinc-aluminum-magnesium coating or aluminum-zinc coating.
  • Table 1 lists the mass percentages of each chemical element in the steel plate substrate of the coated ultra-high-strength steel of Examples 1-28 and the comparative steel of Comparative Examples 1-4.
  • Table 2-1 and Table 2-2 list the specific processes in the above process steps (1)-(6) for the coated ultra-high-strength steel of Examples 1-28 and the comparative steel of Comparative Examples 1-4. parameter.
  • Table 3 lists the coated ultra-high strength steel substrates of Examples 1-28 and the comparative steel substrates of Comparative Examples 1-4. Observation results of the microstructure of the material. Detection was carried out according to GB/T 15749-2008 quantitative metallographic determination method.
  • Tensile property test GB/T228.1-2010 Metal Materials
  • Tensile Test Part 1 Room Temperature Test Method was used to conduct a tensile test to test the coated ultra-high-strength steel base materials of Examples 1-28 and Comparative Example 1. Comparative yield strength, tensile strength, uniform elongation and elongation at break of -4 steel base materials.
  • Table 4 lists the observed results of mechanical properties of the coated ultra-high strength steel substrates of Examples 1-28 and the comparative steel substrates of Comparative Examples 1-4.
  • Table 5 lists the specific spot welding process parameters of the coated ultra-high strength steel of Examples 1-28 and the comparative steel plates of Comparative Examples 1-4.
  • TSS tensile shear
  • CTS cross tensile specimen
  • metallographic specimen welded under each welding current.
  • TSS and CTS can be correspondingly pressed using a tensile testing machine.
  • ISO 14273-2016 and ISO 14272-2016 standards measured the joint bearing capacity of the finished steel plate samples of each example and comparative example after spot welding. The results are shown in Table 6 below.
  • the obtained metallographic samples first use dilute hydrochloric acid to remove the coating on the joint surface, and observe the distribution and direction of spot welding cracks under a microscope. Select the section that passes through the center of the nugget and can cut the most spot welding cracks as the The metallographic section of the joint was sampled using wire cutting. The section includes all the welding characteristic areas of the spot welding joint. The surface of the intercepted sample was rinsed to prevent foreign matter from interfering with the test results. The rinsed sample was dried; The dried samples were mounted, ground and polished, and measured under a metallographic microscope. The maximum crack length was filled in Table 6.
  • Table 6 lists the mechanical properties and solder joint crack test results of the coated ultra-high-strength steel of Examples 1-28 and Comparative Examples 1-4 after spot welding.
  • Table 6 The two values of the nominal coating weight in Table 6 refer to the coating weight on the front and back of the steel plate.
  • GI means pure zinc coating
  • GA It means zinc-iron alloy coating
  • ZM means zinc-aluminum-magnesium coating
  • AZ aluminum-zinc coating.
  • the microstructure of the substrates of the steel materials of Examples 1-28 and Comparative Examples 1-4 designed by the present invention are both: ferrite + martensite + retained austenite.
  • the microstructure of the substrates of each embodiment and comparative example meets the following indicators: the volume proportion of ferrite is 25% to 45%, and the volume proportion of martensite is 45% to 65%; among them, the size is less than 10 ⁇ m
  • the volume proportion of ferrite grains is ⁇ 85%, and the volume proportion of ferrite grains with a size below 5 ⁇ m is ⁇ 55%;
  • the average grain size of retained austenite is ⁇ 2 ⁇ m, and the volume proportion of retained austenite is ⁇ 2 ⁇ m.
  • the average C content in the body is ⁇ 1.0%.
  • the steel base plates of Examples 1-28 and Comparative Examples 1-4 designed by the invention all have very excellent mechanical properties, and their mechanical properties all meet the design requirements of the invention, and the Examples
  • the yield strength of the coated ultra-high-strength steel substrate of 1-28 is specifically between 755 and 845MPa, its tensile strength is between 1040 and 1140MPa, its uniform elongation is between 14.7-16.4%, and its fracture elongation is between Between 22.6-25.4%.

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Abstract

本发明公开了一种低点焊裂纹敏感性的带镀层超高强钢，其包括钢基板以及钢基板表面的含锌镀层，所述钢基板含有Fe和不可避免的杂质元素，所述钢基板还含有质量百分含量如下的下述各化学元素：C：0.10～0.30％；Si：0.30～0.90％；Mn：1.00～2.20％；B：0.001～0.003％；Al：0.30～1.00％。相应地，本发明还公开了上述带镀层超高强钢的制造方法，采用该制造方法可以有效制备本发明上述的带镀层超高强钢，其最终的屈服强度为600MPa～850MPa，抗拉强度为980MPa～1150MPa，均匀延伸率不小于13％，断裂延伸率不小于15％。

Description

一种低点焊裂纹敏感性的带镀层超高强钢及其制造方法 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
目前，随着汽车轻量化进程的推进，以及汽车行业对车辆耐腐蚀能力要求的提高，众多车企对带镀层的超高强度钢需求也变得越来越大，对于冷冲压钢来说，镀层主要是含锌镀层，如热镀纯锌、热镀锌铁及电镀纯锌等镀层。电阻点焊因其焊接速度快、成本低等优势，成为汽车焊接生产中最主要的连接方法。很多含锌镀层超高强钢点焊时接头常常表现出较好的准静态力学性能，但在与电极接触的钢板表面及其附近区域出现一种焊接裂纹，并将这种焊接裂纹称为点焊裂纹，图1就示意性地显示了带锌镀层的超高强钢的点焊裂纹的示意图。然而，抑制这种裂纹形成的难度非常高，因此，这种裂纹已成为含锌镀层超高强钢在汽车领域应用的重要障碍及瓶颈。
现有技术中已知一些技术方案对焊接裂纹进行控制，但仍然存在诸多缺点：
例如：公开号为CN108015401A，公开日为2018年5月11日，名称为“具有良好接头性能的镀锌高强钢电阻点焊方法”的中国专利文献，公开了一种通过点焊工艺创新，在保证焊点接头性能不降低的同时抑制点焊表面裂纹的方法。
又例如：公开号为CN109385515A，公开日为2019年2月26日，名称为“多层钢以及降低液态金属脆化的方法”的中国专利文献，公开了一种抑制高强钢点焊裂纹产生的方法，其将高强钢先脱碳再镀锌的方式抑制点焊裂纹的产生，以将脱碳层厚度在10至50微米之间。
再例如：公开号为CN110892087A，公开日为2020年3月17日，名称为“具有高可电阻点焊性的经锌涂覆的钢板”的中国专利文献，公开了一种通过 提高露点，使钢板形成内氧化层，以改善表层成分分布，从而降低点焊裂纹发生的技术方案。但是，该技术方案存在着内氧化层厚度和均匀性以及钢板表层成分分布均匀性不易控制的问题。
发明内容
本发明的目的之一在于提供一种低点焊裂纹敏感性的带镀层超高强钢，该带镀层超高强钢具有十分优异的质量和性能，其在满足用户对带镀层高强钢性能及点焊接头力学性能要求的同时，还具有低点焊裂纹敏感性，其具有良好的推广应用前景。
为了实现上述目的，本发明提供了一种低点焊裂纹敏感性的带镀层超高强钢，其包括钢基板以及钢基板表面的含锌镀层，所述钢基板含有Fe和不可避免的杂质元素，所述钢基板还含有质量百分含量如下的下述各化学元素：
C：0.10～0.30％；Si：0.30～0.90％；Mn：1.00～2.20％；B：0.001～0.003％；Al：0.30～1.00％。
进一步地，在本发明所述的带镀层超高强钢中，所述钢基板的各化学元素质量百分含量为：
C：0.10～0.30％；Si：0.30～0.90％；Mn：1.00～2.20％；B：0.001～0.003％；Al：0.30～1.00％；余量为Fe和不可避免的杂质元素。
进一步地，在本发明所述的带镀层超高强钢中，所述钢基板的各化学元素还含有Mo：0.10～2.00％。
在本发明上述的技术方案中，其采用了以碳、硅、锰、硼或硼钼复合为主的化学成分设计，充分利用了碳、硅、锰、硼、钼元素在材料相变中的作用，从而实现了本发明所述的超高强钢在高力学性能与高点焊性能的统一，进而最终获得具有低点焊裂纹敏感性的带镀层超高强钢产品。
在本发明所述的带镀层超高强钢中，各化学元素的设计原理如下所述：
C：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，碳在奥氏体中的溶解度远高于其在铁素体中的溶解度，其可以延长奥氏体转变前的孕育期，降低Ms温度。钢中碳的质量百分比越高，则残余奥氏体的分数越多，配分时碳在残余奥氏体中的富集程度越高，有利于增强残余奥氏体稳定性，产生TRIP效应，提高材料延展性。此外，碳也是钢中最基本的固溶强化元素。但 需要注意的是，钢中碳含量也不宜过高，当碳的质量百分比过高时，会降低钢材的焊接性，尤其会显著增大接头的焊点表面裂纹发生倾向。基于此，为了发挥C元素的有益效果，在本发明中，将C元素的质量百分含量控制在0.10～0.30％。
当然，在一些优选的实施方式中，为了获得更优的实施效果，可以进一步优选地将C元素的质量百分含量控制在0.15～0.20％之间。
Si：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，硅在碳化物中的溶解度极小，其能够在配分处理过程中强烈抑制渗碳体形成，促进碳向残余奥氏体中富集，提高残余奥氏体的稳定性。但需要注意的是，钢中不宜添加过量的硅，硅的质量百分比过高会降低钢的高温塑性，并在钢板表面形成稳定氧化物，降低钢板的润湿性，尤其是本案发明人发现硅是显著增大焊点表面裂纹倾向的一种元素。基于此，在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，具体将Si元素的质量百分含量控制在0.30～0.90％。
当然，在一些优选的实施方式中，为了获得更优的实施效果，可以进一步优选地将Si元素的质量百分含量控制在0.40～0.80％之间。
Mn：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，锰能扩大奥氏体相区，降低Ac3、Ms和Mf点，提高奥氏体稳定性和钢的淬透性，降低临界转变速率，其有利于将残余奥氏体保存至室温，同时锰在钢中也能起到固溶强化效果，当钢中Mn元素含量过低时，会因偏析导致低冷速下产生铁素体、珠光体带状组织。另外，钢中锰元素的质量百分含量也不宜过高，当钢中锰的质量百分含量过高时，会加剧晶粒粗化趋势，降低钢的塑性和韧性，并恶化耐腐蚀性能，尤其是会增加镀层下母材浅表层锰的富集，增大接头的焊点表面裂纹敏感性，恶化焊接性能。基于此，为了发挥Mn元素的有益效果，在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，将Mn元素的质量百分含量控制在1.00～2.20％。
当然，在一些优选的实施方式中，为了获得更优的实施效果，可以进一步优选地将Mn元素的质量百分含量控制在1.00～2.00％之间。
B：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，硼能显著提高钢的淬透性。一方面，硼易在晶界偏聚，填充了晶界缺陷，降低了晶界能，使原本是铁素体形核地的奥氏体晶界上新相形核困难增加，奥氏体稳定性增 强，从而提高了淬透性；另一方面，硼元素的添加使钢对碳、锰等增加淬透性元素的需求减少。
另外，本案发明人还发现硼元素会降低锰元素向晶界的偏聚，从而降低了锰在镀层下钢基体的富集层形成，从而显著降低点焊裂纹的形成。但硼的含量也不是越多越好，当晶界缺陷被填满后，若仍有更多的硼非平衡偏聚，则会在晶界形成“硼相”沉淀，增加晶界能，同时“硼相”将作为新相的核心，促使形核速度增加，使奥氏体稳定性下降，使淬透性下降。即有明显的“硼相”析出对淬透性有不良影响，并且大量的“硼相”析出会使钢变脆，给钢的力学性能带来不好的影响；且脆化的晶界会促进点焊裂纹的发生及扩展。
本案发明人研究发现，0.0030％的硼含量是使晶界发生脆化的拐点。基于此，为了发挥B元素的优异效果，在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，将B元素的质量百分含量控制在0.001～0.003％。
当然，在一些优选的实施方式中，为了获得更优的实施效果，可以进一步优选地将B元素的质量百分含量控制在0.0015～0.003％之间。
Al：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，铝以固溶态存在时，能增加层错能，抑制渗碳体析出和γ到马氏体转变，提高奥氏体稳定性。并且铝与碳、氮形成细小弥散分布的难溶质点，可以细化晶粒，但是铝的强化效果弱于硅，其稳定奥氏体的能力也较硅弱。此外，当钢中铝的质量百分含量过高时，容易形成大量氧化物夹杂，不利于炼钢连铸。因此，在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，将Al元素的质量百分含量控制在0.30～1.00％之间。
Mo：在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，钼元素使C曲线右移的同时，能够使Ms点下降，从而提高钢的淬透性，且在提高钢强度的同时提高钢的延伸率；另外，钼是强碳化物形成元素，其在冶炼时会形成细小且弥散分布的MoC颗粒，且会在后续相变时以硬质第二相的形式分布在马氏体中，从而显提高钢的韧性。更重要的是，本案发明人发现，钼的添加能显著提高钢材的抗点焊裂纹的能力。基于此，综合考虑上述因素及钼的添加成本，在本发明所述的低点焊裂纹敏感性的带镀层超高强钢中，将Mo元素的质量百分含量控制在0.10～2.00％之间。在一些实施方案中，将Mo元素的质量百分含量控制在0.15～2.00％之间。
当然，在一些优选的实施方式中，为了获得更优的实施效果，可以进一步优选地将Mo元素的质量百分含量控制在0.10～1.00％之间。在一些实施方案中，将Mo元素的质量百分含量控制在0.15～1.00％之间或在0.20～1.00％之间。
进一步地，在本发明所述的带镀层超高强钢中，所述钢基板的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：0.15～0.20％；
Si：0.40～0.80％；
Mn：1.00～2.00％；
B：0.0015～0.003％；
Mo：0.10～1.00％。
进一步地，在本发明所述的带镀层超高强钢中，所述钢基板的不可避免的杂质元素的质量百分含量满足：P≤0.01％，S≤0.005％。
在上述技术方案中，P、S元素均是本发明所述的带镀层超高强钢中的杂质元素，在技术条件允许情况下，为了获得性能更好且质量更优的钢材，应尽可能降低带镀层超高强钢中杂质元素的含量。
在本发明中，P、S元素均为杂质元素，其中，P虽然能起到固溶强化作用，抑制碳化物形成，有利于提高残余奥氏体的稳定性，但是P的质量百分比过高会弱化晶界，增大材料脆性，恶化焊接性能，也就是说P元素的正面作用弱于其负面作用，因此，优选地将P的质量百分比控制为P≤0.01％。
相应地，钢中S元素易在晶界处形成低熔共晶，其质量百分含量过高时，将会明显恶化材料的塑性，因此将S元素的质量百分含量控制为S≤0.005％。
进一步地，在本发明所述的带镀层超高强钢中，所述钢基板的微观组织为铁素体+马氏体+残余奥氏体。
进一步地，在本发明所述的带镀层超高强钢中，其中铁素体的体积比例为25％～45％；并且/或者马氏体的体积比例为45％～65％。
进一步地，在本发明所述的带镀层超高强钢中，其中在铁素体中，尺寸为10μm以下的晶粒的体积占比≥85％，尺寸为5μm以下的晶粒的体积占比≥55％。
进一步地，在本发明所述的带镀层超高强钢中，其中残余奥氏体的平均晶粒尺寸≤2μm；并且/或者残余奥氏体中的平均C含量≥1.0％，如1.0-1.3％。
进一步地，在本发明所述的带镀层超高强钢中，所述“低点焊裂纹敏感性”指：采用点焊工艺进行焊接时不产生焊点表面裂纹，或者，若产生了焊点表面裂纹，则焊点表面裂纹的最大长度小于板厚的5％，优选小于板厚的3％，更优选小于板厚的1％。
进一步地，在本发明所述的带镀层超高强钢中，其力学性能满足：其屈服强度≥600MPa，如≥700MPa，优选≥750MPa；抗拉强度≥980MPa，优选≥1040MPa；均匀延伸率≥13％，优选≥15.5％；断裂延伸率≥15％，优选≥20％，更优选≥22.5％。
进一步地，在本发明所述的带镀层超高强钢中，其力学性能满足：其屈服强度为600MPa～850MPa，抗拉强度为980MPa～1150MPa，均匀延伸率不小于13％，断裂延伸率不小于15％。
进一步地，在本发明所述的带镀层超高强钢中，其力学性能满足：其屈服强度为755～845MPa，抗拉强度为1040～1140MPa，均匀延伸率为14.7～16.4％，断裂延伸率在22.6～25.4％之间。
进一步地，在本发明所述的带镀层超高强钢中，其镀层为纯锌镀层、锌铁合金镀层、锌铝镁镀层或铝锌镀层。
进一步地，在本发明所述的带镀层超高强钢中，镀层重量为30～120g/m2，优选为50～100g/m2。
相应地，本发明的另一目的在于提供一种带镀层超高强钢的制造方法，该制造方法生产简单，所得到的高强度钢在同等力学性能条件下，点焊性能尤其是抗焊点表面裂纹的能力得到显著提高。
为了实现上述目的，本发明提出了一种带镀层超高强钢的制造方法，其具体包括步骤：
(1)冶炼和薄板连铸；
(2)加热；
(3)热轧：控制热轧后钢带表面氧化皮厚度≤4μm，并且热轧后带钢表面氧化皮中的FeO+Fe3O4的质量百分含量≤50wt％；
(4)酸洗，或者酸洗+冷轧；
(5)连续退火：在800～920℃退火，然后以3～10℃/s的冷速缓冷至700～770℃；再以50～500℃/s(如100～500℃/s、200～500℃/s或400～500℃/s)的 冷速快冷至200～300℃；然后再加热至360～460℃，保温50～600s；最后冷却至室温；
(6)镀覆含锌镀层。
在本发明所述的技术方案中，该制造方法通过采用薄板坯连铸工艺配合酸洗或酸轧工艺，并在连续退火、镀层制造后能够获得低点焊裂纹敏感性的带镀层超高强钢。此制造方法生产的带镀层超高强钢可以使用车厂常规的点焊工艺进行焊接，且其焊点表面裂纹的最大长度小于板厚的5％，其具有很低的点焊裂纹敏感性，具有十分优异的质量。
在本发明上述步骤(1)中，由于采用的是薄板坯连铸，因此可以省却粗轧工序，减小热轧变形量，从而保证在后续的步骤(4)以及步骤(5)的钢板性能。此外，由于步骤(1)中采用的是薄板坯连铸，其可以充分利用板坯热量，降低加热所需能耗，从而获得更加均匀的铁素体或铁素体+珠光体组织，进而有利于步骤(6)中成品的基板微观组织中保持一定量的细晶粒铁素体，提高组织均匀性。
而在步骤(2)中，控制热轧后钢带表面氧化皮厚度≤4μm，并且热轧后带钢表面氧化皮中的(FeO+Fe3O4)≤50wt％，有利于后续在步骤(4)的进行，并对连续退火后所获得的钢板性能有着重要影响，这是因为：在本发明所述的技术方案中，FeO、Fe3O4比Fe2O3更加难酸洗，而控制本发明所制备的热轧后带钢表面氧化皮厚度以及热轧后带钢表面氧化皮中的(FeO+Fe3O4)≤50wt％，可以有效提高酸洗效果，获得可用于直接连续退火的酸洗板表面，而由于酸洗板可以直接进行连续退火，使得热轧组织变形量小，钢板组织以珠光体与铁素体为主；因而，在相同连续退火条件下可以降低材料强度，使得组织更加均匀，从而获得优异的延展性。
相应地，在步骤(5)中，发明人针对连续退火工艺进行了优化设计，其通过控制800～920℃退火温度进行退火，可以形成均匀化的奥氏体组织或奥氏体+铁素体组织；然后以3～10℃/s的冷速缓冷至700～770℃，以进一步调整组织中铁素体的含量，获得一定比例的铁素体，从而提高材料的塑形；之后以再以50～500℃/s的速度冷却至200～300℃(即介于Ms(马氏体转变的起始温度)与Mf(马氏体转变终了温度)之间)，此时，奥氏体部分转变为马氏体，其能够保证钢材获得较高的强度；之后，再加热至360～460℃并保温50～600s， 能够使碳在马氏体和奥氏体中发生配分，形成一定量富碳的残余奥氏体，稳定保持到室温，由于TRIP效应，可以显著提高钢的加工硬化能力和成形性，得到延展性优异的高强度钢板。
另外，鉴于本发明的发明人对碳、硅、锰、硼或硼钼对点焊裂纹影响的认识，尤其是硅、锰易在钢板浅表层富集显著增大焊点表面裂纹敏感性的认知，本发明除了在成分设计时对钢材的碳、硅、锰、硼或硼钼的含量进行了严格限定，其与同强度级别钢种相比，还减少了碳、硅、锰元素的含量，并复合添加了硼和钼元素来增加钢的淬透性，从而能够确保在步骤(5)中获得所需的微观组织成分，并保证了通过步骤(6)后获得的含锌镀层的超高强钢具有低点焊裂纹敏感性，即焊点表面裂纹的最大长度值小于板厚的5％。
需要说明的是，在步骤(6)中镀覆含锌镀层时，可以但不限于采用热镀、电镀、真空镀技术生产获得含锌镀层。
综上所述，由于本发明所设计的这种带镀层超高强钢采用的是：碳、硅、锰、硼、铝、钼的化学成分设计并配合进行了铁素体晶粒细化。因而，在连续退火过程中，在奥氏体逆相变的形核点增多的同时，能够进一步细化晶粒尺寸，从而稳定保持到室温的残余奥氏体的平均晶粒尺寸≤2μm；残余奥氏体中的平均C含量≥1.0％。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(1)中，控制薄带连铸出口端的板坯厚度为50～58mm。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(1)中，控制薄带连铸的拉速为2～5m/min。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(2)中，将板坯加热到1200～1250℃。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(3)中，控制终轧温度为860～930℃，卷取温度为450～600℃。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(4)中，当采用酸洗+冷轧时，控制冷轧变形量为40％～60％。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，在步骤(5)中，控制连续退火炉内的还原性气氛中氢气的体积含量为10～15％。
进一步地，在本发明所述的带镀层超高强钢的制造方法中，步骤(5)的 退火工艺参数满足下述各项的至少其中之一：
退火温度为820～870℃；
以3～10℃/s的冷速缓冷至700～730℃；
快冷至250～300℃；
快冷后再加热至400～430℃，保温180～300s。
相较于现有技术，本发明所述的低点焊裂纹敏感性的带镀层超高强钢及其制造方法具有如下所述的优点以及有益效果：
在本发明所设计的这种带镀层超高强钢中，其化学元素成分以C、Si、Mn、B、Al和Mo为基础，通过优化各元素的配比，可以获得具有低点焊裂纹敏感性的带镀层超高强钢。
在本发明所述的低点焊裂纹敏感性的带镀层超高强钢具有十分优异的质量和性能，其在满足用户对带镀层高强钢性能及点焊接头力学性能要求的同时，还具有低点焊裂纹敏感性。
采用本发明这种技术方案所制备的带镀层超高强钢包括钢基板以及钢基板表面的含锌镀层，其力学性能满足：其屈服强度为600MPa～850MPa，抗拉强度为980MPa～1150MPa，均匀延伸率不小于13％，断裂延伸率不小于15％；并且，在实际采用点焊工艺进行焊接时，若产生了焊点表面裂纹，则焊点表面裂纹的最大长度小于板厚的5％。
本发明所设计的制造方法的生产工艺简单，所获的高强度钢在同等力学性能条件下，抗点焊裂纹能力显著提高，在下游用户安全结构件生产中将具有较好的应用前景。
附图说明
图1示意性地显示了带锌镀层的超高强钢的点焊裂纹的示意图。
具体实施方式
下面将结合具体的实施例对本发明所述的低点焊裂纹敏感性的带镀层超高强钢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1-28和对比例1-4
本发明所述的实施例1-28的带镀层超高强钢和对比例1-4的对比钢材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薄板连铸，其中，控制薄带连铸出口端的板坯厚度为50～58mm，控制薄带连铸的拉速为2～5m/min，以获得对应的板坯。
(2)加热：将板坯加热到1200～1250℃。
(3)热轧：控制热轧后钢带表面氧化皮厚度≤4μm，并且热轧后带钢表面氧化皮中的(FeO+Fe3O4)≤50wt％，控制终轧温度为860～930℃，卷取温度为450～600℃。
(4)酸洗，或者酸洗+冷轧，当采用酸洗+冷轧时，控制其冷轧变形量为40％～60％，当酸洗后直接镀锌或锌合金时，表2-2中的冷轧变形量为0。
(5)连续退火：控制连续退火炉内的还原性气氛中氢气的体积含量为10～15％；在800～920℃退火，优选地可以控制退火温度为820～870℃，然后以3～10℃/s的冷速缓冷至700～770℃，以获得一定比例的铁素体，缓冷终冷温度可以优选地控制在700～730℃之间；而后再以50～500℃/s的冷速快冷至200～300℃，以使奥氏体部分转变为马氏体，快冷终冷温度可以优选地控制在250～300℃之间；然后快冷后再加热至360～460℃，优选为400～430℃，保温50～600s，优选保温时间为180～300s；最后冷却至室温。
(6)镀覆锌或锌合金镀层，其镀层可以具体选用为纯锌镀层、锌铁合金镀层、锌铝镁镀层或铝锌镀层。
需要注意的是，在本发明中，实施例1-28的带镀层超高强钢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1-4的对比钢材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1-28的带镀层超高强钢和对比例1-4的对比钢材的钢板基板的各化学元素的质量百分配比。
表1.(wt％，余量为Fe和除P、S以外其他不可避免的杂质)

相应地，表2-1和表2-2列出了实施例1-28的带镀层超高强钢和对比例1-4的对比钢材在上述工艺步骤(1)-(6)中的具体工艺参数。
表2-1.


注：在上述表2-1中，“(FeO+Fe3O4)占比”表示热轧后带钢表面氧化皮中的FeO和Fe3O4占氧化皮
的质量百分比。
表2-2.

需要注意的都是，在进行上述步骤(6)的镀覆工艺前，为了验证实施例1-28的带镀层超高强钢具有十分优异的性能，发明人在进行镀覆工艺前，将经过步骤(5)的连续退火工艺后所获得的实施例1-28的带镀层超高强钢基材和对比例1-4的对比钢材基材分别进行了取样，并对各实施例和对比例的基材进行了微观组织观察，同时对各实施例和对比例的基材的力学性能进行了检测。其中，各实施例和对比例的基材的微观组织观察结果列于下述表3之中，各实施例和对比例的基材的力学性能检测结果列于下述表4之中。
表3列出了实施例1-28的带镀层超高强钢基材和对比例1-4的对比钢材基 材的微观组织观察结果。根据GB/T 15749-2008定量金相测定方法进行检测。
表3.
相应地，在针对各实施例和对比例基材进行力学性能检测时，相关力学性能检测手段如下所述：
拉伸性能测试：采用GB/T228.1-2010金属材料拉伸试验第1部分：室温试验方法进行拉伸试验，以测试获得实施例1-28的带镀层超高强钢基材和对比例1-4的对比钢材基材的屈服强度、抗拉强度、均匀延伸率和断裂延伸率。
表4列出了实施例1-28的带镀层超高强钢基材和对比例1-4的对比钢材基材的力学性能观察结果。
表4.
相应地，为了验证经过上述步骤(1)-(6)所制备的成品实施例1-28的带镀层超高强钢具有相当优异的低点焊裂纹敏感性，针对最终所制备的成品实施例1-28的带镀层超高强钢和对比例1-4的对比钢板，发明人分别进行取样，并针对各实施例和对比例的成品钢板样品进行了点焊工艺进行焊接试验，相关点焊工艺参数参见下述表5。
表5列出了实施例1-28的带镀层超高强钢和对比例1-4的对比钢板的具体点焊工艺参数。
表5.

在实际进行点焊工艺时，每个焊接电流下分别焊接拉剪(TSS)试样、十字拉伸(CTS)试样、金相试样各一个，TSS、CTS能够对应采用拉伸试验机按ISO 14273-2016及ISO 14272-2016标准测量各实施例和对比例成品钢板样品经点焊焊接的接头承载力，结果见下述表6。
相应地，针对获得的金相试样，先用稀盐酸去除接头表面的镀层，并在显微镜下观测点焊裂纹的分布及走向，选取过熔核中心且可切割最多点焊裂纹的截面做为接头的金相剖切面，采用线切割进行取样，截面包含点焊接头所有焊接特征区域，对截取后的试样进行表面冲洗，防止异物对测试结果的干扰，冲洗后的样品进行干燥处理；对干燥后的试样进行镶样、磨抛处理，并在金相显微镜下进行测量，将最大裂纹长度填写在表6中。
表6列出了实施例1-28的带镀层超高强钢和对比例1-4的在进行点焊后的焊点力学性能及焊点裂纹测试结果。
表6.


注：表6中镀层公称重量的两个数值是指钢板的正面与背面的镀层重量。GI表示纯锌镀层，GA
表示锌铁合金镀层，ZM表示锌铝镁镀层，AZ表示铝锌镀层。
由本发明上述的表3可以看出，本发明所设计的实施例1-28与对比例1-4的钢材的基板的微观组织均为：铁素体+马氏体+残余奥氏体。并且，各实施例和对比例的基板的微观组织均满足以下指标：铁素体的体积比例为25％～45％、马氏体的体积比例为45％～65％；其中，尺寸为10μm以下的铁素体晶粒的体积占比均≥85％，尺寸为5μm以下的铁素体晶粒的体积占比均≥55％；残余奥氏体的平均晶粒尺寸均≤2μm，残余奥氏体中的平均C含量均≥1.0％。
由上述表4可以看出，发明所设计的实施例1-28与对比例1-4的钢材的基板的均具有十分优异的力学性能，其力学性能均满足本发明的设计要求，且实施例1-28的带镀层超高强钢的基板的屈服强度具体在755～845MPa之间，其抗拉强度在1040～1140MPa之间，其均匀延伸率在14.7-16.4％之间，其断裂延伸率在22.6-25.4％之间。
同时由表6可知，对于本案实施例1-28的成品带镀层超高强钢，当焊接电流小于飞溅电流时焊接的焊点无点焊裂纹；当焊接电流大于飞溅电流时，焊接的焊点最长的裂纹也小于板厚的5％；而对于本案所设计的对比例1-4的对比例钢板，不论焊接电流是否大于飞溅电流，所焊接的焊点都会有裂纹，且飞溅焊点的裂纹比无飞溅焊点的裂纹严重，且最大裂纹长度与板厚之比均远大于 5％。由此，说明了采用本发明技术方案所设计的实施例1-28的各实施例的带镀层超高强钢在保证板材性能的同时具有优异的低点焊裂纹敏感性。
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 (21)

  1. 一种低点焊裂纹敏感性的带镀层超高强钢，其包括钢基板以及钢基板表面的含锌镀层，所述钢基板含有Fe和不可避免的杂质元素，其特征在于，所述钢基板还含有质量百分含量如下的下述各化学元素：
    C：0.10～0.30％；Si：0.30～0.90％；Mn：1.00～2.20％；B：0.001～0.003％；Al：0.30～1.00％。
  2. 如权利要求1所述的带镀层超高强钢，其特征在于，所述钢基板的各化学元素质量百分含量为：
    C：0.10～0.30％；Si：0.30～0.90％；Mn：1.00～2.20％；B：0.001～0.003％；Al：0.30～1.00％；余量为Fe和不可避免的杂质元素。
  3. 如权利要求1或2所述的带镀层超高强钢，其特征在于，所述钢基板的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：0.15～0.20％；
    Si：0.40～0.80％；
    Mn：1.00～2.00％；
    B：0.0015～0.003％。
  4. 如权利要求1或2所述的带镀层超高强钢，其特征在于，所述钢基板的化学元素还含有Mo：0.10～2.00％；优选地，Mo的质量百分含量为0.15～2.00％。
  5. 如权利要求4所述的带镀层超高强钢，其特征在于，所述钢基板的Mo质量百分含量为0.10～1.00％；优选地，Mo的质量百分含量为0.15～1.00％。
  6. 如权利要求1或2所述的带镀层超高强钢，其特征在于，所述钢基板的不可避免的杂质元素的质量百分含量满足：P≤0.01％，S≤0.005％。
  7. 如权利要求1或2所述的带镀层超高强钢，其特征在于，所述钢基板的微观组织为铁素体+马氏体+残余奥氏体。
  8. 如权利要求7所述的带镀层超高强钢，其特征在于，其中铁素体的体积比例为25％～45％；并且/或者马氏体的体积比例为45％～65％。
  9. 如权利要求7所述的带镀层超高强钢，其特征在于，其中在铁素体中，尺 寸为10μm以下的晶粒的体积占比≥85％，尺寸为5μm以下的晶粒的体积占比≥55％。
  10. 如权利要求7所述的带镀层超高强钢，其特征在于，其中残余奥氏体的平均晶粒尺寸≤2μm；并且/或者残余奥氏体中的平均C含量≥1.0％。
  11. 如权利要求1或2所述的带镀层超高强钢，其特征在于，采用点焊工艺进行焊接时，若产生了焊点表面裂纹，则焊点表面裂纹的最大长度小于板厚的5％。
  12. 如权利要求1或2所述的带镀层超高强钢，其特征在于，其力学性能满足：其屈服强度为600MPa～850MPa，抗拉强度为980MPa～1150MPa，均匀延伸率不小于13％，断裂延伸率不小于15％。
  13. 如权利要求1或2所述的带镀层超高强钢，其特征在于，其镀层为纯锌镀层、锌铁合金镀层、锌铝镁镀层或铝锌镀层。
  14. 如权利要求1-13中任意一项所述的带镀层超高强钢的制造方法，其特征在于，包括步骤：
    (1)冶炼和薄板连铸；
    (2)加热；
    (3)热轧：控制热轧后钢带表面氧化皮厚度≤4μm，并且热轧后带钢表面氧化皮中的FeO+Fe3O4的质量百分含量≤50wt％；
    (4)酸洗，或者酸洗+冷轧；
    (5)连续退火：在800～920℃退火，然后以3～10℃/s的冷速缓冷至700～770℃；再以50～500℃/s的冷速快冷至200～300℃；然后再加热至360～460℃，保温50～600s；最后冷却至室温；
    (6)镀覆含锌镀层。
  15. 如权利要求14中所述的制造方法，其特征在于，在步骤(1)中，控制薄带连铸出口端的板坯厚度为50～58mm。
  16. 如权利要求14中所述的制造方法，其特征在于，在步骤(1)中，控制薄带连铸的拉速为2～5m/min。
  17. 如权利要求14中所述的制造方法，其特征在于，在步骤(2)中，将板坯加热到1200～1250℃。
  18. 如权利要求14中所述的制造方法，其特征在于，在步骤(3)中，控制终 轧温度为860～930℃，卷取温度为450～600℃。
  19. 如权利要求14中所述的制造方法，其特征在于，在步骤(4)中，当采用酸洗+冷轧时，控制冷轧变形量为40％～60％。
  20. 如权利要求14中所述的制造方法，其特征在于，在步骤(5)中，控制连续退火炉内的还原性气氛中氢气的体积含量为10～15％。
  21. 如权利要求14中所述的制造方法，其特征在于，步骤(5)的退火工艺参数满足下述各项的至少其中之一：
    退火温度为820～870℃；
    以3～10℃/s的冷速缓冷至700～730℃；
    快冷至250～300℃；
    快冷后再加热至400～430℃，保温180～300s。
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