WO2022206917A1 - 高成形性热镀铝锌或热镀锌铝镁双相钢及其快速热处理热镀制造方法 - Google Patents

高成形性热镀铝锌或热镀锌铝镁双相钢及其快速热处理热镀制造方法 Download PDF

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WO2022206917A1
WO2022206917A1 PCT/CN2022/084543 CN2022084543W WO2022206917A1 WO 2022206917 A1 WO2022206917 A1 WO 2022206917A1 CN 2022084543 W CN2022084543 W CN 2022084543W WO 2022206917 A1 WO2022206917 A1 WO 2022206917A1
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hot
aluminum
magnesium
dip galvanized
dip
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PCT/CN2022/084543
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English (en)
French (fr)
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王健
李俊
张理扬
杜小峰
丁志龙
刘华飞
任玉苓
杜瑶
林传华
杨奕
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宝山钢铁股份有限公司
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Priority claimed from CN202110360129.3A external-priority patent/CN115181840A/zh
Priority claimed from CN202110360134.4A external-priority patent/CN115181885B/zh
Application filed by 宝山钢铁股份有限公司 filed Critical 宝山钢铁股份有限公司
Priority to JP2023560592A priority Critical patent/JP2024512730A/ja
Priority to KR1020237037740A priority patent/KR20230166117A/ko
Priority to EP22779097.9A priority patent/EP4317513A1/en
Priority to US18/552,934 priority patent/US20240167140A1/en
Publication of WO2022206917A1 publication Critical patent/WO2022206917A1/zh

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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
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    • 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
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • 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
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    • 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
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    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
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Definitions

  • the invention belongs to the technical field of rapid heat treatment of materials, in particular to high formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel (including hot-dip galvanized aluminum-zinc AZ products and hot-dip galvanized aluminum-magnesium AM products) and rapid heat treatment hot-dip galvanized steels thereof Manufacturing method.
  • hot-dip galvanized dual-phase steel is the most widely used and has the best application prospect.
  • Low-carbon and low-alloy hot-dip galvanized dual-phase steel has the characteristics of low yield point, high initial work hardening rate, and good matching of strength and plasticity. It has become a widely used steel for automotive structural stamping with high strength and good formability.
  • hot-dip pure zinc has gradually been unable to meet the requirements, and it is urgent to develop new varieties of high corrosion resistance coatings. Therefore, there are more and more studies on hot-dip galvanized aluminum-zinc and hot-dip galvanized aluminum-magnesium coatings with better corrosion resistance. Correspondingly, hot-dip galvanized, hot-dip galvanized aluminum-magnesium high-strength steel products also came into being.
  • the main method for the development of hot-dip galvanized aluminum-magnesium-coated dual-phase steel is to change the microstructure and properties of hot-dip galvanized dual-phase steel by adding alloying elements and adjusting the soaking temperature, time and cooling rate in the critical annealing process. .
  • Chinese patent application CN201710994660.X discloses "Hot-dip Al-Zn steel sheet for 550MPa grade structure and its preparation method", the chemical composition is C: 0.02-0.07%, Si ⁇ 0.03%, Mn: 0.15-0.30%, P ⁇ 0.020% , Si ⁇ 0.020%, Nb: 0.015-0.030%, Als: 0.020-0.070%, cold rolling with a low cold rolling reduction rate of 55-60%, yield strength above 550MPa, tensile strength of 560MPa, elongation
  • the steel plate proposed in this patent has the problem of low elongation and high yield strength, which will affect the subsequent processing.
  • Chinese patent application CN102363857B has published "a production method of color-coated sheet for structure with yield strength of 550MPa", in which Ti and Nb are at most 0.05% and 0.045%, respectively, and the yield strength Rp 0.2 reaches 550-600MPa; tensile strength Rm It is 560-610MPa, and the elongation after fracture is A 80 mm ⁇ 6%.
  • the strengthening method mainly maintains most of the unrecrystallized band-like structure through low-temperature annealing to improve the strength, but the plasticity is poor, which also affects the forming.
  • Chinese patent application CN100529141C discloses "a full-hard aluminum-zinc-coated steel sheet and its production method", the yield strength of the steel sheet prepared by this method reaches more than 600MPa, the elongation at break is less than or equal to 7%, and the total content of Ti and Nb is 0.15%- 0.100%, the annealing temperature is controlled at 630-710 ° C, and the full hard steel plate is obtained by low-temperature recovery annealing. The elongation of the steel plate product obtained by this method is too low to meet the current processing requirements for formability.
  • Chinese patent application CN201911161556.8 discloses "a hot-dip galvanized aluminum-magnesium high-strength steel, preparation method and application".
  • the hot-dip galvanized aluminum-magnesium high-strength steel includes a base material and a zinc-aluminum-magnesium alloy coating on the surface of the base material.
  • Design, as well as the control of the production process on the basis of the composition design, have formed the CSP thin slab continuous casting and rolling production line and the ordinary hot-dip galvanizing production line as the core process of smelting, hot rolling, cold rolling, and annealing. Process production plan and core production technology.
  • the yield strength of the above hot-dip galvanized aluminum-magnesium high-strength steel is greater than 550MPa, and the elongation is greater than 17%. Due to poor formability, it is only suitable for industries such as photovoltaic brackets and highway guardrails that require high corrosion resistance but low formability requirements.
  • Chinese patent application CN104419867A published "a 1250MPa grade ultra-high strength zinc-aluminum-magnesium coated steel sheet and its production method", the chemical composition weight percentage of the steel sheet is: C: 0.15-0.35%, Si: 0.50-1.80%, Mn: 2.0- 5.0%, Mn/Si not less than 2, the rest are iron and inevitable impurities; the chemical composition weight percentage of the coating is: Al: 1-15%, Mg: 1-5%, Al/Mg ⁇ 1, and the rest is Zn and inevitable impurities.
  • the production method includes smelting-continuous casting-hot rolling-cold rolling-continuous hot-dip plating process.
  • the high corrosion-resistant and ultra-high-strength zinc-aluminum-magnesium-coated steel sheet manufactured according to the present invention has a strength of 1250-1500 MPa and an elongation after fracture of 12 ⁇ 18%, the corrosion resistance is more than 4 times that of ordinary galvanized sheet, and the coating has no cracks or peeling when bent at 180°5a, which meets the requirements of high corrosion resistance and high strength reduction.
  • this patent proposes a kind of production method of high-strength zinc-aluminum-magnesium coated steel sheet, the excessive Si content is prone to surface quality problems, and the C content is too high, and the weldability is poor, which affects subsequent processing and forming.
  • High corrosion-resistant advanced high-strength steels represented by dual-phase steels have broad application prospects, and rapid heat treatment technology has great development value. The combination of the two will surely provide more space for the development of dual-phase steels.
  • the purpose of the present invention is to provide a high formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel (including hot-dip galvanized AZ and hot-dip galvanized aluminum-magnesium AM products) with a tensile strength of ⁇ 590 MPa, and a rapid heat treatment thermal
  • the plating manufacturing method controls the recovery of the deformed matrix, ferrite recrystallization, austenite transformation and grain growth during the annealing process by rapid heating, and obtains fine ferrite structure and polymorphism after the final heat treatment.
  • the strength of the material is greatly improved and the toughness is also improved.
  • the yield strength of the dual-phase steel is ⁇ 300MPa, the tensile strength is ⁇ 590MPa, the elongation is ⁇ 20%, the strength-plastic product is ⁇ 15GPa%, and the strain hardening index
  • the n90 value is greater than 0.20.
  • the technical scheme of the present invention is:
  • High formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel with tensile strength ⁇ 590MPa its chemical composition mass percentage is: C: 0.045-0.12%, Si: 0.1-0.5%, Mn: 1.0-2.0% , P ⁇ 0.02%, S ⁇ 0.006%, Al: 0.02 ⁇ 0.055%, can also contain one or two of Cr, Mo, Ti, Nb, V, and Cr+Mo+Ti+Nb+V ⁇ 0.5 %, the balance is Fe and other inevitable impurities.
  • the high formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel has a yield strength ⁇ 300 MPa, a tensile strength ⁇ 590 MPa, an elongation ⁇ 20%, a strong-plastic product ⁇ 15GPa%, and a strain hardening index n 90 value is greater than 0.20; more preferably, the yield strength is 300-560MPa, such as 300-400MPa or 450-560MPa, the tensile strength is 590-860MPa, preferably 630-860MPa, the elongation is 20-30%, and the strong-plastic product is 15 ⁇ 21 GPa%.
  • the C content is 0.045-0.105% or 0.05-0.12%.
  • the Si content is 0.1-0.4%.
  • the Mn content is 1.0%-1.5% or 1.2-2.0%.
  • the dual-phase steel may contain one or two of Cr, Mo, Ti, Nb, and V, and Cr+Mo+Ti+Nb+V ⁇ 0.3%.
  • the metallographic structure of the dual-phase steel is a uniformly distributed ferrite and martensite dual-phase structure, and the average grain size is 1-5 ⁇ m.
  • the high formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel with tensile strength ⁇ 590 MPa is obtained by the following process:
  • the cold rolling reduction rate is 40-85%;
  • the cold-rolled steel plate is rapidly heated to 750-845°C, and the rapid heating adopts a one-stage or two-stage type;
  • the heating rate is 15-500°C/s (eg 50-500°C/s);
  • the first stage is heated from room temperature to 550-650°C at a heating rate of 15-500°C/s, and the second stage is heated at a heating rate of 30-500°C/s (such as 50-500°C/s). Heating rate from 550 ⁇ 650°C to 750 ⁇ 845°C;
  • soaking temperature 750 ⁇ 845°C
  • soaking time 10 ⁇ 60s
  • hot-dip aluminum-zinc dipping After hot-dip aluminum-zinc dipping, rapidly cool to room temperature at a cooling rate of 30-200°C/s (such as 30-150°C/s) to obtain hot-dip aluminum-zinc AZ products; or,
  • hot-dip galvanizing aluminum-magnesium After hot-dip galvanizing aluminum-magnesium, it is rapidly cooled to room temperature at a cooling rate of 10-300°C/s (eg, 30-180°C/s) to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the C content is 0.065-0.085% or 0.07-0.10%.
  • the Si content is 0.15-0.25% or 0.1-0.4%.
  • the Mn content is 1.2%-1.35% or 1.5-1.8%.
  • the dual-phase steel may contain one or two of Cr, Mo, Ti, Nb, and V, and Cr+Mo+Ti+Nb+V ⁇ 0.4%, or ⁇ 0.2%.
  • the whole process of rapid heat treatment and hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium in step 4) takes 29-159s, preferably 29-122s.
  • the hot rolling finishing temperature is ⁇ A r3 .
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the heating rate is 50-300°C/s.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 15-300°C/s, and the second stage is heated at a heating rate of 50-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 30-300°C/s, and the second stage is heated at a heating rate of 80-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • step 4 after hot-dip galvanizing aluminum-magnesium, it is rapidly cooled to room temperature at a cooling rate of 30-250° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the metallographic structure of the dual-phase steel of the present invention is a uniformly distributed ferrite and martensite dual-phase structure, and the average grain size is 1-5 ⁇ m, such as 1-3 ⁇ m.
  • the chemical composition mass percentage of the high formability hot-dip galvanized aluminum-zinc or hot-dip galvanized aluminum-magnesium dual-phase steel with a tensile strength ⁇ 590 MPa of the present invention is: C: 0.045-0.105%, Si: 0.1- 0.4%, Mn: 1.0 ⁇ 1.5%, P ⁇ 0.02%, S ⁇ 0.006%, Al: 0.02 ⁇ 0.055%, and can also contain one or two of Cr, Mo, Ti, Nb, V, and Cr+ Mo+Ti+Nb+V ⁇ 0.3%, the balance is Fe and other inevitable impurities.
  • the C content of the dual-phase steel is 0.065-0.085%; preferably, the Si content of the dual-phase steel is 0.15-0.25%; preferably, the Mn content of the dual-phase steel is 1.2%-1.35%.
  • one or two of Cr, Mo, Ti, Nb, and V may be contained in the dual-phase steel, and Cr+Mo+Ti+Nb+V ⁇ 0.2%.
  • the metallographic structure of the dual-phase steel is a uniformly distributed ferrite and martensite dual-phase structure, and the average grain size is 1-3 ⁇ m.
  • the dual-phase steel has a yield strength of 30 to 400 MPa, a tensile strength of 630 to 700 MPa, an elongation of 22 to 30%, a strong-plastic product of 15 to 20 GPa%, and a strain hardening index n 90 value greater than 0.21; more preferably
  • the yield strength of this dual-phase steel is 304-398MPa, the tensile strength is 630-698MPa, the elongation is 22.3-29.4%, the strength-plastic product is 15.3-19.4GPa%, and the strain hardening index n90 value is greater than 0.21.
  • this dual phase steel is obtained by the following process:
  • the cold rolling reduction rate is 40-85%;
  • the cold-rolled steel plate is rapidly heated to 750-845°C, and the rapid heating adopts a one-stage or two-stage type;
  • the heating rate is 15-500°C/s (eg 50-500°C/s);
  • the first stage is heated from room temperature to 550-650°C at a heating rate of 15-500°C/s, and the second stage is heated at a heating rate of 30-500°C/s (such as 50-500°C/s). Heating rate from 550 ⁇ 650°C to 750 ⁇ 845°C;
  • soaking temperature 750 ⁇ 845°C
  • soaking time 10 ⁇ 60s
  • the cooling rate of /s (such as 50 ⁇ 150°C/s) is rapidly cooled to 580 ⁇ 600°C, immersed in a zinc pot for hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium;
  • hot-dip galvanizing aluminum-magnesium After hot-dip galvanizing aluminum-magnesium, it is cooled to room temperature at a cooling rate of 30-180° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the whole process of the rapid heat treatment and hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium in step d) takes 29-122 s.
  • the hot rolling finishing temperature is ⁇ A r3 .
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the heating rate is 50-300°C/s.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 15-300°C/s, and the second stage is heated at a heating rate of 50-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 30-300°C/s, and the second stage is heated at a heating rate of 80-300°C/s
  • the heating rate is from 550 to 650 °C to 750 to 845 °C.
  • the chemical composition mass percentage of the high formability hot-dip galvanized aluminum-magnesium dual-phase steel with tensile strength ⁇ 590MPa is: C: 0.05-0.12%, Si: 0.1-0.5%, Mn: 1.2 ⁇ 2.0%, P ⁇ 0.015%, S ⁇ 0.003%, Al: 0.02 ⁇ 0.055%, and can also contain one or two of Cr, Mo, Ti, Nb, V, and Cr+Mo+Ti +Nb+V ⁇ 0.5%, the balance is Fe and other inevitable impurities.
  • the C content of the dual-phase steel is 0.07-0.10%.
  • the Si content of the dual-phase steel is 0.1-0.4%.
  • the Mn content of the dual-phase steel is 1.5-1.8%.
  • the dual-phase steel may contain one or two of Cr, Mo, Ti, Nb, and V, and Cr+Mo+Ti+Nb+V ⁇ 0.4%.
  • the metallographic structure of the dual-phase steel is a uniformly distributed ferrite and martensite dual-phase structure, and the average grain size is 1-5 ⁇ m.
  • the dual-phase steel has a yield strength of 470-560 MPa, a tensile strength of 780-860 MPa, an elongation of 20-25%, a strong-plastic product of 16-21 GPa%, and a strain hardening index n 90 value greater than 0.20; more preferably The yield strength of this dual-phase steel is 476-556MPa, the tensile strength is 786-852MPa, the elongation is 20.1-24.8%, the strength-plastic product is 16.7-20.2GPa%, and the strain hardening index n90 value is greater than 0.20.
  • this dual phase steel is obtained by the following process:
  • the cold rolling reduction rate is 40-85%;
  • the cold-rolled steel plate is rapidly heated to 750-845°C, and the rapid heating adopts a one-stage or two-stage type;
  • the heating rate is 15-500°C/s (eg 50-500°C/s);
  • the first stage is heated from room temperature to 550-650°C at a heating rate of 15-500°C/s, and the second stage is heated at a heating rate of 30-500°C/s (such as 50-500°C/s). Heating rate from 550 ⁇ 650°C to 750 ⁇ 845°C;
  • soaking temperature 750 ⁇ 845°C
  • soaking time 10 ⁇ 60s
  • hot-dip aluminum-zinc dipping After hot-dip aluminum-zinc dipping, rapidly cool to room temperature at a cooling rate of 30-150°C/s to obtain hot-dip aluminum-zinc AZ products; or,
  • hot-dip galvanizing aluminum-magnesium After hot-dip galvanizing aluminum-magnesium, it is rapidly cooled to room temperature at a cooling rate of 10-300° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the whole process of rapid heat treatment and hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium in step D) takes 29-159s.
  • the hot rolling finishing temperature is ⁇ A r3 .
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • step D) when the rapid heating adopts one-stage heating, the heating rate is 50-300°C/s.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 15-300°C/s, and the second stage is heated at a heating rate of 50-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-650°C at a heating rate of 30-300°C/s, and the second stage is heated at a heating rate of 80-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • step D) after hot-dip galvanizing aluminum-magnesium, it is rapidly cooled to room temperature at a cooling rate of 30-250° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • Carbon is the most common strengthening element in steel. Carbon increases the strength of steel and reduces its plasticity. However, for forming steel, low yield strength, high uniform elongation and total elongation are required. Therefore, carbon The content should not be too high. Carbon content has a great influence on the mechanical properties of steel. With the increase of carbon content, the number of pearlite will increase, and the strength and hardness of steel will be greatly improved, but its plasticity and toughness will decrease significantly. high, there will be obvious network carbides in the steel, and the existence of the network carbides will significantly reduce the strength, plasticity and toughness, and the strengthening effect caused by the increase of the carbon content in the steel will also be significantly weakened. , so that the process performance of the steel deteriorates, so the carbon content should be reduced as much as possible under the premise of ensuring the strength.
  • carbon mainly affects the volume fraction of austenite formed during annealing.
  • austenite the diffusion process of carbon in austenite or ferrite actually plays a role in Controls the process of austenite grain growth.
  • the volume fraction of austenite increases, and then the martensite phase structure formed after cooling increases, and the strength of the material increases. Therefore, the strength and toughness matching of the material and the strength of the rapid annealing process are considered comprehensively. improvement.
  • the carbon content is limited within the range of 0.045-0.12%.
  • Mn Manganese can form a solid solution with iron, thereby improving the strength and hardness of ferrite and austenite in carbon steel, and enabling the steel to obtain finer and higher strength pearlite in the cooling process after hot rolling, and pearlite.
  • the content of Mn also increases with the increase of Mn content.
  • Manganese is also a carbide forming element, and manganese carbides can dissolve into cementite, thereby indirectly enhancing the strength of pearlite.
  • Manganese can also strongly enhance the hardenability of steel, further increasing its strength.
  • manganese is one of the elements that significantly affects the kinetics of austenite formation during critical zone annealing.
  • Manganese mainly affects the transformation and growth of austenite into ferrite after the formation of Final equilibration process with ferrite. Since the diffusion rate of manganese in austenite is much smaller than that in ferrite, the austenite grains controlled by manganese diffusion take a longer time to grow, while manganese reaches uniformity in austenite. The distribution time will be longer. When heating in the critical zone, if the holding time is short, the manganese element cannot be uniformly distributed in the austenite, and then the cooling rate is insufficient, and the uniform martensite island structure will not be obtained.
  • the manganese content is generally high, resulting in a high manganese content after the formation of austenite, which ensures the hardenability of the austenite island.
  • the manganese element expands the ⁇ phase region and reduces the temperature of A c1 and A c3 , so the manganese-containing steel will obtain a higher martensite volume fraction than the low-carbon steel under the same heat treatment conditions.
  • the present invention designs the manganese content within the range of 1.0-2.0%.
  • Si Silicon forms a solid solution in ferrite or austenite, thereby enhancing the yield strength and tensile strength of steel, and silicon can increase the cold working deformation hardening rate of steel, which is a beneficial element in alloy steel.
  • silicon is obviously enriched on the intergranular fracture surface of silicon-manganese steel. The segregation of silicon at the grain boundary can slow down the distribution of carbon and phosphorus along the grain boundary, thereby improving the embrittlement state of the grain boundary.
  • Silicon can improve the strength, hardness and wear resistance of steel without significantly reducing the plasticity of steel. Silicon has strong deoxidation ability and is a commonly used deoxidizer in steelmaking. Silicon can also increase the fluidity of molten steel, so silicon is generally contained in steel, but when the content of silicon in steel is too high, its plasticity and toughness will decrease significantly.
  • the main effect of silicon is to reduce the volume fraction of austenite at final equilibrium for a given annealing time. Silicon has no obvious effect on the growth rate of austenite, but has a significant effect on the formation and distribution of austenite. Therefore, the present invention determines the silicon content within the range of 0.1-0.5%.
  • Chromium and iron form a continuous solid solution, reducing the austenite phase region. Chromium forms various carbides with carbon, and its affinity with carbon is greater than that of iron and manganese. Chromium and iron can form intermetallic compound ⁇ phase (FeCr). Chromium reduces the concentration of carbon in pearlite and the limit solubility of carbon in austenite. Chromium slows down the decomposition rate of austenite and significantly improves the hardenability of steel. But it also increases the temper brittleness tendency of steel.
  • Chromium element can improve the strength and hardness of steel while adding other alloying elements, the effect is more significant. Since Cr improves the quenching ability of steel during air cooling, it has an adverse effect on the weldability of steel. However, when the chromium content is less than 0.3%, the adverse effects on weldability can be ignored; when the content is greater than this, defects such as cracks and slag inclusions are likely to occur during welding. When Cr coexists with other alloying elements (such as coexisting with V), the adverse effect of Cr on weldability is greatly reduced. For example, when Cr, Mo, V and other elements exist in the steel at the same time, even if the Cr content reaches 1.7%, there is no significant adverse effect on the welding performance of the steel. In the present invention, chromium element is a beneficial and unnecessary addition element, and the addition amount should not be too much considering factors such as cost increase. In some embodiments, the Cr content is ⁇ 0.35%.
  • Molybdenum suppresses the self-diffusion of iron and the diffusion rate of other elements.
  • the atomic radius of Mo is larger than that of ⁇ -Fe atoms.
  • Mo can increase the bond attraction of lattice atoms and increase the recrystallization temperature of ⁇ ferrite.
  • the strengthening effect of Mo in pearlitic, ferritic, martensitic steels, and even in high-alloy austenitic steels is also very obvious.
  • the good effect of Mo in steel also depends on the interaction with other alloying elements in the steel.
  • the solid solution strengthening effect of Mo is more significant. This is because when the strong carbide forming element is combined with C to form a stable carbide, it can promote the more effective dissolution of Mo into the solid solution, which is more conducive to the improvement of the thermal strength of the steel. Adding Mo can also increase the hardenability of steel, but the effect is not as significant as that of C and Cr. Mo will inhibit the transformation of the pearlite region and accelerate the transformation of the medium temperature region, so the steel containing Mo can also form a certain amount of bainite and eliminate the formation of ferrite when the cooling rate is large. One of the reasons why the thermal strength of alloy heat-resistant steel has a favorable effect.
  • Mo can also significantly reduce the hot brittle tendency of steel and reduce the speed of pearlite spheroidization.
  • Mo content is below 0.15%, there is no adverse effect on the weldability of the steel.
  • molybdenum element is a beneficial and unnecessary addition element, and the addition amount should not be too much considering factors such as cost increase.
  • the Mo content is ⁇ 0.2%
  • Micro-alloying elements Ti, Nb, V adding trace amounts of micro-alloying elements Nb, V, Ti to the steel can ensure that the steel can pass through its carbon and nitride dots (size less than 5nm) when the carbon equivalent is low. Dispersion and precipitation and solid solution of Nb, V, Ti, refine the grains, greatly improve the strength and toughness of the steel, especially the low temperature toughness, so that the steel has good weldability and usability.
  • Nb, V, Ti are carbide and nitride forming elements. These elements can meet this requirement at relatively low concentrations.
  • Nb, V, Ti are strong carbide forming elements.
  • Nb, V and Ti can prevent the growth of austenite grains and increase the roughening temperature of the steel. This is because their carbon and nitride dispersed small particles can fix the austenite grain boundaries, hinder the migration of austenite grain boundaries, increase the recrystallization temperature of austenite, and expand the unrecrystallized area, that is, Austenite grain growth is prevented.
  • Adding trace amounts of Nb, V, and Ti to steel can, on the one hand, increase the strength while reducing the carbon equivalent content, thereby improving the weldability of the steel; on the other hand, it can fix impure substances such as oxygen, nitrogen, sulfur, etc. , thereby improving the weldability of the steel; secondly, due to the effect of its microscopic particles, such as the insolubility of TiN at high temperatures, it can prevent the coarsening of the grains in the heat-affected zone and improve the toughness of the heat-affected zone, thereby improving the steel. Welding performance.
  • the microalloying elements are beneficial and unnecessary added elements, and the added amount should not be too much considering factors such as cost increase.
  • the Ti content is ⁇ 0.04%.
  • the Nb content is ⁇ 0.05%.
  • the V content is ⁇ 0.05%.
  • the invention controls the recovery, recrystallization, austenite transformation and grain growth of the deformed structure in the continuous heat treatment process through the rapid heat treatment process of rapid heating, short-term heat preservation and rapid cooling, and not only forms ferrite in the cooling process
  • the matrix phase is formed, and various strengthening phases and the composition gradient distribution in the phases are produced, and finally a fine ferrite structure and a multi-morphic strengthening phase structure are obtained, so that the material can obtain a better combination of strength and toughness, and reduce the cost of alloying and various processes It is difficult to manufacture and improve the welding performance of steel of the same strength grade.
  • the specific principle is that different heating rates are used in different temperature stages of the heating process.
  • the recovery of the deformed structure mainly occurs in the low temperature section, and a relatively low heating rate can be used to reduce energy consumption; the recrystallization and grain growth of different phase structures mainly occur in the high temperature section.
  • a relatively high heating rate must be used to shorten the residence time of the structure in the high temperature range to ensure that the grain cannot grow.
  • the high-formability hot-dip galvanized aluminum-zinc and hot-dip galvanized aluminum-magnesium dual-phase steel with a tensile strength of ⁇ 590 MPa comprises the following steps:
  • the cold rolling reduction rate is 40 to 85%, and the rolled hard strip or steel plate is obtained after cold rolling;
  • the heating rate is 15-500°C/s (eg 50-500°C/s);
  • the first stage is heated from room temperature to 550-650°C at a heating rate of 15-500°C/s, and the second stage is heated at a heating rate of 30-500°C/s (such as 50-500°C/s). Heating rate from 550 ⁇ 650°C to 750 ⁇ 845°C;
  • Soaking is carried out at the target temperature of 750-845 °C in the two-phase region of austenite and ferrite, and the soaking time is 10-60s;
  • the strip or steel plate After soaking, the strip or steel plate is cooled slowly to 670-770°C at a cooling rate of 5-15°C/s; then rapidly cooled to 580-600°C at a cooling rate of 50-200°C/s (eg 50-150°C/s). °C;
  • hot-dip galvanizing aluminum-magnesium After hot-dip galvanizing aluminum-magnesium, it is rapidly cooled to room temperature at a cooling rate of 10-300°C/s (30-180°C/s) to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the whole process of the rapid heat treatment of hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium is 29-159 s, such as 29-122 s.
  • the method is used to prepare the dual-phase steel according to any of the foregoing embodiments whose chemical composition mass percentage is as follows: C: 0.045-0.105%, Si: 0.1-0.4%, Mn : 1.0 ⁇ 1.5%, P ⁇ 0.02%, S ⁇ 0.006%, Al: 0.02 ⁇ 0.055%, can also contain one or two of Cr, Mo, Ti, Nb, V, and Cr+Mo+Ti+ Nb+V ⁇ 0.3%, the balance is Fe and other unavoidable impurities; wherein, after hot-dip galvanizing, cool down to room temperature at a cooling rate of 30-200°C/s to obtain hot-dip AZ products, or, After hot-dip galvanizing aluminum-magnesium, it is cooled to room temperature at a cooling rate of 30-180° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the method is used to prepare the dual-phase steel according to any of the foregoing embodiments whose chemical composition mass percentage is as follows: C: 0.05-0.12%, Si: 0.1-0.5%, Mn: 1.2 ⁇ 2.0%, P ⁇ 0.015%, S ⁇ 0.003%, Al: 0.02 ⁇ 0.055%, and can also contain one or two of Cr, Mo, Ti, Nb, V, and Cr+Mo+Ti +Nb+V ⁇ 0.5%, the balance is Fe and other unavoidable impurities; among them, after hot-dip galvanizing, it is rapidly cooled to room temperature at a cooling rate of 30-150°C/s to obtain hot-dip AZ products, Alternatively, after hot-dip galvanizing aluminum-magnesium, rapidly cool to room temperature at a cooling rate of 10-300° C./s to obtain a hot-dip galvanizing aluminum-magnesium AM product.
  • the whole process of rapid heat treatment and hot-dip galvanizing or hot-dip galvanizing aluminum-magnesium is as follows: C: 0.05-0.1
  • the hot rolling finishing temperature is ⁇ A r3 .
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the heating rate is 50-300°C/s.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-650°C at a heating rate of 15-300°C/s, and the second stage is heated at a heating rate of 50-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-650°C at a heating rate of 30-300°C/s, and the second stage is heated at a heating rate of 80-300°C/s
  • the heating rate is from 550 ⁇ 650°C to 750 ⁇ 845°C.
  • the final temperature of the rapid heating is 770-830°C.
  • step 4 After the strip or steel sheet is heated to the target temperature in the austenite and ferrite two-phase region, soaking is performed while keeping the temperature unchanged.
  • the strip or steel plate is heated or cooled in a small range during the soaking time period, and the temperature after heating is not more than 845°C, and the temperature after cooling is not lower than 750°C.
  • the soaking time is 10-40s.
  • the recrystallization kinetics of the continuous heating process can be quantitatively described by the relationship affected by the heating rate.
  • the functional relationship between the ferrite recrystallization volume fraction and the temperature T during the continuous heating process is:
  • X(t) is the volume fraction of ferrite recrystallization
  • n is the Avrami index, which is related to the phase transformation mechanism and depends on the decay cycle of the recrystallization nucleation rate, generally in the range of 1 to 4
  • T is Heat treatment temperature
  • T star is the recrystallization start temperature
  • ⁇ is the heating rate
  • b(T) is obtained by the following formula:
  • the traditional heat treatment process adopts slow heating. Under this condition, the deformed matrix undergoes recovery, recrystallization and grain growth in sequence, and then the phase transformation from ferrite to austenite occurs.
  • the phase deformation nucleus is mainly concentrated in the iron that has grown. At the grain boundary of the element body, the nucleation rate is low, and the final grain structure is relatively coarse.
  • the phase transformation of ferrite to austenite occurs before the recovery of the deformed matrix is completed, or the austenite transformation occurs just after the recrystallization is completed, and the grains have not grown up.
  • the grain size is small and the grain boundary area is large, the nucleation rate is significantly increased, and the austenite grains are obviously refined.
  • a large number of nucleation points are provided for austenite due to the retention of a large number of crystal defects such as dislocations in the ferrite crystal, which makes austenite appear Explosive nucleation, further refinement of austenite grains.
  • the retained high-density dislocation line defects also become the channels for the high-speed diffusion of carbon atoms, so that each austenite grain can rapidly grow and grow, thus increasing the austenite volume fraction.
  • the present invention sets the heating rate as 15-500°C/s during one-stage rapid heating, and sets the heating rate as 2-stage rapid heating as 30 ⁇ 500°C/s. .
  • the optimal heating rate in different heating temperature ranges are also different: from 20°C to 550-650°C, the heating rate has the greatest impact on the recovery process, and the heating rate is controlled to be 15-500°C/s, more preferably 30-500°C/s; the heating temperature is from 550-650°C When the austenitizing temperature is 750-845°C, the heating rate has the greatest influence on the grain growth process, and the heating rate is controlled to be 50-500°C/s; more preferably, it is 80-500°C/s.
  • the soaking temperature usually depends on the C content.
  • the C content in the dual-phase steel of the present invention is 0.045-0.12%, and the A C1 and A C3 of the present steel are about 730°C and 870°C, respectively.
  • the rapid heat treatment process of the present invention the strip steel is heated to between A C1 and A C3 for soaking, and the rapid heating technology is used to retain a large number of dislocations in the insufficiently recrystallized ferrite, which provides a good solution for austenite transformation. Larger nucleation driving force, so compared with the traditional continuous annealing process, the rapid heat treatment method of the present invention can obtain more and finer austenite structure.
  • the present invention takes the lead in proposing that soaking temperature is increased and decreased within a certain range: namely, soaking zone temperature is ramped up and ramped down, but soaking temperature must be kept within a certain range.
  • the advantage of this is that the rapid heating and cooling process in the temperature range of the two-phase region is actually to further increase the degree of superheat and subcooling, which is convenient for the rapid phase transition process.
  • the grains can be further refined through repeated ferrite to austenite phase transformation and austenite to ferrite phase transformation, and at the same time, it also has a certain influence on the formation of carbides and the uniform distribution of alloying elements, and finally forms Smaller structure and uniform distribution of alloying elements.
  • a small amount of fine granular undissolved carbides should be evenly distributed in the steel, which can not only prevent the abnormal growth of austenite grains, but also increase the content of each alloying element in the matrix accordingly, so as to improve the The purpose of mechanical properties such as strength and toughness of alloy steel.
  • the selection of soaking temperature should also aim to obtain fine and uniform austenite grains, avoid coarse austenite grains, and achieve the purpose of obtaining fine martensite structure after cooling. Too high soaking temperature will make the austenite grains coarse, and the martensite structure obtained after rapid cooling will also be coarser, resulting in poor mechanical properties of the steel; it will also increase the amount of retained austenite and reduce martensite. The number of bodies reduces the hardness and wear resistance of the steel. Too low soaking temperature will cause insufficient carbon and alloying elements dissolved in austenite, resulting in uneven distribution of alloying element concentration in austenite, greatly reducing the hardenability of steel, and adversely affecting the mechanical properties of steel. influences.
  • the soaking temperature of hypoeutectoid steel should be Ac3+30 ⁇ 50°C.
  • the existence of carbide-forming elements will hinder the transformation of carbides, so the soaking temperature can be appropriately increased.
  • the present invention selects 750-845°C as the soaking temperature, in order to obtain a more ideal and more reasonable final structure.
  • the influencing factors of soaking time also depend on the content of carbon and alloying elements in the steel.
  • the content increases, not only will the thermal conductivity of the steel decrease, but also because the diffusion rate of alloying elements is slower than that of carbon elements, alloying elements will The structural transformation of the steel is obviously delayed, and the holding time should be appropriately extended at this time. Since this process adopts rapid heating, the material contains a large number of residual dislocations in the two-phase region, so a large number of nucleation points are provided for the formation of austenite, and a fast diffusion channel for carbon atoms is provided, so austenite can be extremely fast.
  • the present invention sets the holding time as 10-60s.
  • the control of the rapid cooling process needs to combine the comprehensive factors such as the evolution results of each structure and the diffusion distribution of alloy elements in the pre-heating and soaking process to ensure that the ideal structure of each phase and the material structure with reasonable distribution of elements are finally obtained.
  • the cooling rate of the sample during quenching must be greater than the critical cooling rate to obtain the martensite structure, and the critical cooling rate mainly depends on the material composition, and the Si content in the present invention is 0.1-0.5 %, the Mn content is 1.0-2.0%, and the content is relatively high, so Si and Mn greatly enhance the hardenability of the dual-phase steel, thereby reducing the critical cooling rate.
  • the cooling rate also needs to comprehensively consider the microstructure evolution and alloy diffusion distribution results during the heating process and soaking process, so as to finally obtain a reasonable microstructure distribution of each phase and alloy element distribution.
  • the rapid cooling rate is set to 50-150°C/s.
  • the rapid heat treatment process reduces the residence time of the strip in the high-temperature furnace, so the enrichment of alloy elements on the surface of the high-strength strip during the heat treatment process
  • the amount of galvanized steel is significantly reduced, which is beneficial to the improvement of the platability of high-strength hot-dip galvanized products, and thus to the reduction of surface leakage defects and the improvement of corrosion resistance of high-strength hot-dip galvanized products, thereby improving the yield;
  • the refinement of the grains and the reduction of the alloy content of the material make the dual-phase steel products obtained by the technology of the present invention also improve the processing and forming properties such as the hole expansion performance and the bending performance, and the user performance such as the welding performance.
  • the rapid heat treatment technology reduces the heating process and soaking process time, shortens the length of the furnace (at least one third shorter than the traditional continuous annealing furnace), and significantly reduces the number of furnace rolls, which reduces the probability of surface defects in the furnace, so the product Surface quality will be significantly improved.
  • the present invention has the following advantages:
  • the present invention suppresses the recovery of the deformed structure and the ferrite recrystallization process during the heat treatment process through rapid heat treatment, so that the recrystallization process overlaps with the austenite transformation process, and the shape of the recrystallized grains and austenite grains is increased.
  • Nucleation point, shorten the grain growth time the metallographic structure of the obtained dual-phase steel is a uniformly distributed ferrite and martensite dual-phase structure, and the fine martensite in the structure after rapid heat treatment is characterized by having Block, strip, granular and other forms, and the distribution is more uniform, so that dual-phase steel products can obtain good strong-plastic matching.
  • the average grain size of the dual-phase steel obtained by the rapid heat treatment of the present invention is 1 to 5 ⁇ m, a good effect of grain refinement strengthening can be obtained.
  • the yield strength is 300-560MPa
  • the tensile strength is 590-860MPa
  • the elongation is 20-30%
  • the strength-plastic product is 15-21GPa%
  • the strain hardening index n 90 value is greater than 0.20.
  • the time of the whole heat treatment process can be shortened to 29-159s, which greatly reduces the time of the whole heat treatment process.
  • the traditional continuous annealing process time is usually 300-480s), which significantly improves production efficiency, reduces energy consumption, and reduces production costs.
  • the rapid heat treatment method of the present invention shortens the length of the heating section and soaking section of the continuous hot-dip galvanizing annealing furnace (compared to the traditional continuous
  • the annealing furnace can shorten at least one third), time and the entire heat treatment process time, which can save energy, reduce emissions, reduce consumption, and significantly reduce the one-time investment in furnace equipment, significantly reducing production and operation costs and equipment maintenance costs; Products with strength grades can reduce the alloy content, reduce the production cost of heat treatment and pre-process, and reduce the manufacturing difficulty of each process before heat treatment.
  • the enrichment of alloying elements on the surface of high-strength strip steel is significantly reduced, which is conducive to the improvement of product platability, and thus is conducive to the reduction of leakage plating defects and resistance to the surface of high-strength hot-dip galvanized aluminum-zinc and hot-dip galvanized aluminum-magnesium products.
  • the dual-phase steel products obtained by the technology of the present invention have process forming properties such as hole expansion properties and bending properties, and welding properties. User performance such as performance has also improved.
  • the hot-dip galvanized aluminum-zinc and hot-dip galvanized aluminum-magnesium dual-phase steel obtained by the present invention is used by users such as forming, welding, painting, and corrosion resistance.
  • the performance has also been improved, in which the corrosion resistance is 3 to 8 times that of the traditional GI dual-phase steel.
  • the high formability hot-dip galvanized dual-phase steel obtained by the present invention is beneficial to the development of a new generation of lightweight vehicles, trains, ships, airplanes and other means of transportation, the healthy development of the corresponding industries, and the healthy development of advanced manufacturing industries. are of great value.
  • Fig. 1 is a microstructure picture of a hot-dip galvanized aluminum-magnesium dual-phase steel (AM) produced according to Example 1 (one-stage heating) of test steel A in Example 1 of the present invention.
  • AM hot-dip galvanized aluminum-magnesium dual-phase steel
  • Example 2 is a microstructure picture of the hot-dip Al-Zn dual-phase steel (AZ) produced by Example 1 (two-stage heating) of Test Steel A according to Example 1 of the present invention.
  • Example 3 is a microstructure picture of a hot-dip hot-dip Al-Zn dual-phase steel (AZ) produced by the traditional process 1 of the test steel A in Example 1 of the present invention.
  • Fig. 4 is the microstructure picture of the hot-dip aluminum-zinc dual-phase steel (AZ) produced according to Example 3 (two-stage heating) of Test Steel I in Example 1 of the present invention.
  • Example 5 is a microstructure picture of a hot-dip galvanized aluminum-magnesium dual-phase steel (AM) produced by Example 1 (two-stage heating) of Example 1 of the present invention.
  • AM hot-dip galvanized aluminum-magnesium dual-phase steel
  • FIG. 6 is a microstructure picture of the hot-dip aluminum-zinc dual-phase steel (AZ) produced according to Example 4 (one-stage heating) of Test Steel D in Example 2 of the present invention.
  • FIG. 7 is a microstructure picture of the hot-dip Al-Zn dual-phase steel (AZ) produced by the traditional process 4 of the test steel D in the second embodiment of the present invention.
  • FIG. 8 is a microstructure picture of the hot-dip aluminized zinc dual-phase steel (AZ) produced according to Example 15 (two-stage heating) of the second test steel N of the present invention.
  • FIG. 9 is a microstructure picture of the hot-dip galvanized aluminum-magnesium dual-phase steel (AM) produced according to Example 17 (two-stage heating) of the second test steel E of the present invention.
  • n 90 is carried out in accordance with "GB/T228.1-2010 Tensile Test of Metallic Materials Part 1: Test Method at Room Temperature", using P7 sample to test in the transverse direction, according to "GBT 5028-2008 Tensile Sheet and Strip of Metallic Materials” Determination method of strain hardening index (n value)" to obtain n 90 value.
  • Table 4 is the main performance of the dual-phase steel obtained by the one-stage heating example and the traditional process for the components of the test steel of this embodiment, and the obtained dual-phase steel is prepared by the two-stage heating example and the traditional process for the components of the test steel in Table 5. The main properties of phase steel.
  • the alloy content in the same grade of steel can be reduced, the grains can be refined, and the material structure and the matching of strength and toughness can be obtained.
  • the yield strength of the dual-phase steel obtained by the method of the invention is 304-398 MPa, the tensile strength is 630-698 MPa, the elongation is 22.3-29.4%, the strong-plastic product is 15.3-19.4GPa%, and the strain hardening index n 90 value is greater than 0.21 .
  • FIG. 1 is the microstructure picture of the hot-dip galvanized aluminum-magnesium dual-phase steel (AM) produced by the test steel A of the present invention according to the embodiment 1 (one-stage heating) of the present invention
  • Fig. 2 and Fig. 3 are the typical composition A steel after The structure diagram (two-stage heating) of Example 1 and the comparative traditional process example 1
  • Fig. 4 is the hot-dip Al-Zn dual-phase steel (AZ) produced by Example 3 (two-stage heating) of the steel of this embodiment I Microstructure picture
  • FIG. 5 is the microstructure picture of hot-dip galvanized aluminum-magnesium dual-phase steel (AM) produced by the steel of Example C under the conventional heating rate (two-stage heating) in Example 15.
  • Table 9 is the main properties of the dual-phase steel obtained by the one-stage heating of the test steel composition of the present embodiment and the obtained dual-phase steel by the traditional process, and the table 10 is prepared by the two-stage heating of the obtained embodiment and the traditional process for the test steel composition of the present embodiment. The main properties of the resulting dual-phase steel.
  • the method of the present invention can reduce the alloy content in the steel of the same grade, refine the grains, and obtain the matching of material structure and strength and toughness.
  • the yield strength of the dual-phase steel obtained by the method of the invention is 476-556MPa
  • the tensile strength is 786-852MPa
  • the elongation is 20.1-24.8%
  • the strong-plastic product is 16.7-20.2GPa%
  • the strain hardening index n90 value is greater than 0.20 .
  • Fig. 6, Fig. 7 are the microstructure diagrams of the hot-dip aluminized zinc dual-phase steel (AZ) of the experimental steel D of the present embodiment through the embodiment 4 and the comparative traditional process example 4 (one-stage heating), and Fig. 8 is the experiment of the present embodiment.
  • the microstructure characteristics of the dual-phase steel obtained by the process of the present invention are very fine and evenly distributed in the matrix, which is very beneficial for improving the strength and plasticity of the material. Therefore, the preparation method of the dual-phase steel of the present invention can refine the crystal grains, so that each phase structure of the material is evenly distributed in the matrix, thereby improving the material structure and improving the material properties.
  • the present invention transforms the traditional continuous annealing unit by adopting the rapid heating and rapid cooling process to realize the rapid heat treatment process, which can greatly shorten the length of the heating section and the soaking section of the traditional continuous annealing furnace, and improve the performance of the traditional continuous annealing unit.
  • the unit can realize the advantages of short and compact unit, flexible transition of product specifications and varieties, and strong control ability; for materials, it can refine the grain of the strip steel, further improve the strength of the material, reduce the cost of the alloy and the manufacturing difficulty of the pre-heat treatment process, and improve the material. Forming, welding and other user performance.
  • the invention greatly promotes the technological progress of the continuous annealing process of the cold-rolled steel strip by adopting the rapid heat treatment process. It can be completed within ten seconds or even a few seconds, which greatly shortens the length of the heating section of the continuous annealing furnace, facilitates the improvement of the speed and production efficiency of the continuous annealing unit, and significantly reduces the number of rollers in the furnace of the continuous annealing unit.
  • the number of rollers in the high temperature furnace section of the rapid heat treatment production line is not more than 10, which can significantly improve the surface quality of the strip.
  • the rapid heat treatment process of recrystallization and austenitization completed in a very short time will also provide a more flexible and flexible high-strength steel structure design method, and then without changing the alloy composition and rolling process, etc.
  • the material structure can be improved and the material properties can be improved.
  • Advanced high-strength steel with high corrosion resistance coating represented by dual-phase steel has broad application prospects, and rapid heat treatment technology has great development value. The combination of the two will surely provide greater development and production of dual-phase steel. Space.

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Abstract

抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢及快速热处理热镀制造方法，该钢成分质量百分比为：C：0.045～0.12％，Si：0.1～0.5％，Mn：1.0～2.0％，P≤0.02％，S≤0.006％，Al 0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。本发明通过快速热处理改变退火过程中变形组织的回复、铁素体再结晶及奥氏体相变过程，增加结晶形核点，缩短晶粒长大时间，在提高热处理效率的同时提高了材料的强度和成形性能(n 90值)，扩展了材料性能区间范围。

Description

高成形性热镀铝锌或热镀锌铝镁双相钢及其快速热处理热镀制造方法 技术领域
本发明属于材料快速热处理技术领域，特别涉及高成形性热镀铝锌或热镀锌铝镁双相钢(包括热镀铝锌AZ产品和热镀锌铝镁AM产品)及其快速热处理热镀制造方法。
背景技术
随着人们对能源节约以及材料服役安全意识的逐步提高，很多汽车制造商选择高强钢作为汽车用材，其中汽车的排气系统就要求材料具有高强度高韧性高耐蚀特性，同时还需要一定的耐热特性；同时家电产品及建筑用材不仅对基板材料要高强减薄而且要求涂镀层具有良好的耐蚀性，可见汽车、家电、建筑等领域涂镀产品的耐蚀性、抗凹陷性、耐久强度、大变形冲击强度和安全性要求越来越高。
在汽车用热镀锌高强钢中，热镀锌双相钢应用最为广泛并且应用前景最好。低碳低合金热镀锌双相钢具有屈强点低、初始加工硬化速率高以及强度和塑性匹配性好等特点，成为目前广泛使用的强度高、成形性好的汽车结构冲压用钢。
然而随着对钢制品耐腐蚀性能要求的不断提高，热镀纯锌已经渐渐的不能满足要求，迫切需要开发新的高耐蚀镀层品种。因此，耐腐蚀更优的热镀铝锌、热镀锌铝镁镀层研究越来越多。相应的，热镀铝锌、热镀锌铝镁高强钢产品也应运而生。
目前，针对热镀铝锌、热镀锌铝镁镀层双相钢开发的主要手段是通过添加合金元素、调整临界退火工艺中均热温度、时间和冷速而改变热镀双相钢的组织性能。
中国专利申请CN201710994660.X公开了“550MPa级结构用热镀铝锌钢板及其制备方法”，化学成分为C：0.02-0.07％，Si≤0.03％，Mn：0.15-0.30％，P≤0.020％，Si≤0.020％，Nb：0.015-0.030％，Als：0.020-0.070％，采用55-60％的低冷轧压下率进行冷轧，屈服强度在550MPa以上，抗拉强度为560MPa，伸长率在10％左右，此专利提出的钢板具有伸长率较低，而屈强比较高的问题，会对后续加工过程产生影响。
中国专利申请CN102363857B公布了“一种屈服强度550MPa结构用彩涂板的 生产方法”，其中Ti、Nb最多分别为0.05％和0.045％，其屈服强度Rp 0.2达到550-600MPa；抗拉强度R m为560-610MPa，断后伸长率A 80mm≥6％，强化方式主要通过低温退火保持大部分未再结晶的带状组织，提高强度，但塑性较差，同样对成型产生影响。
中国专利申请CN100529141C公开了“一种全硬质镀铝锌钢板及其生产方法”，该方法制备得到的钢板屈服强度达到600MPa以上，断裂延伸率≤7％，其Ti、Nb总含量0.15％-0.100％，退火温度控制在630-710℃，通过低温回复退火的方式获得全硬质钢板，该方法获得的钢板产品延伸率过低，不能满足目前加工对成型性能要求。
中国专利申请CN201911161556.8公开了“一种热浸镀锌铝镁高强钢、制备方法及应用”，热浸镀锌铝镁高强钢包括基材和基材表面的锌铝镁合金镀层，通过成分设计，以及在成分设计的基础上对于生产工艺过程控制，形成了CSP薄板坯连铸连轧生产线和普通热镀锌生产线为核心工艺的冶炼、热轧、冷轧、退火的工艺生产方案和核心生产技术。以上的热浸镀锌铝镁高强钢的屈服强度大于550MPa，延伸率＞17％。由于成形性较差，只适用于光伏支架、公路护栏等要求高耐蚀但成形性要求不高的行业。
中国专利申请CN106811686A公布了“表面质量好的高强锌铝镁镀层钢板及其制造方法”，钢板的化学成分包含C：0.09-0.18％，Si：0.40-1.60％，Mn：0.80-2.10％，S：0.001-0.008％，还可加入Cr：0.01-0.60％，和/或Mo：0.01-0.30％。镀层的化学成分为Al：1-14％，Mg：1.0-5.0％，其余为锌和不可避免杂质。该专利虽然提出了一种高强的锌铝镁镀层钢板生产方法，但其成本高，而且Si含量过高易出现表面质量问题，屈服强度过高，伸长率较低，影响后续加工和成型。
中国专利申请CN104419867A公布了“一种1250MPa级超高强锌铝镁镀层钢板及其生产方法”，钢板的化学成分重量百分数为：C：0.15～0.35％，Si：0.50～1.80％，Mn：2.0～5.0％，Mn/Si不小于2，其余为铁和不可避免的杂质；镀层的化学成分重量百分数为：Al：1～15％，Mg：1～5％，Al/Mg≥1，其余为Zn和不可避免的杂质。生产方法包括冶炼-连铸-热连轧-冷连轧-连续热浸镀工艺，按照本发明制造的高耐蚀超高强锌铝镁镀层钢板，强度为1250～1500MPa，断后伸长率为12～18％，耐蚀性为普通镀锌板的4倍以上，镀层180°5a弯曲时无裂纹、不剥落，满足高耐蚀高强度减量化需求。该专利虽然提出了一种高强锌铝镁镀层钢板的生产方法，但Si含量过高 易出现表面质量问题，而且C含量过高，焊接性较差，影响后续加工和成型。
综上所述，目前热镀铝锌和热镀锌铝镁产品存在成本高，表面质量较差，强度或伸长率匹配不佳导致后续加工成型的问题。同时，以往受企业生产设备所限，绝大部分的相关研究都是基于现有加热装备的加热速率(5～20℃/s)对带钢进行加热完成再结晶和奥氏体化(中国专利CN104988391A)。近年来，横磁感应加热和新型直火加热等快速加热技术的开发，使快速热处理工艺得以工业化应用。冷轧带钢从室温开始将有可能实现在十几秒甚至几秒内完成奥氏体化过程，大大缩短了炉子加热段长度，提高了机组速度和生产效率。同时，在极短时间内所完成的再结晶和奥氏体化过程也将提供更加灵活及柔性化的组织设计，进而在无需改变合金成分以及轧制工艺的前提下改善材料性能。
以双相钢为代表的高耐蚀先进高强钢有着广阔的应用前景，而快速热处理技术又有着巨大的开发价值，两者的结合必将会为双相钢的开发提供更大的空间。
发明内容
本发明的目的在于提供一种抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢(包括热镀铝锌AZ和热镀锌铝镁AM产品)及快速热处理热镀制造方法，通过快速加热控制退火过程中变形基体的回复、铁素体再结晶、奥氏体相变及晶粒长大等过程，在最后完成热处理后获得细小的铁素体组织及多形态的强化相组织，在材料强度大幅提高的同时韧性亦有所改善，获得双相钢的屈服强度≥300MPa，抗拉强度≥590MPa，延伸率≥20％，强塑积≥15GPa％，应变硬化指数n 90值大于0.20。
为达到上述目的，本发明的技术方案是：
抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢，其化学成分质量百分比为：C：0.045～0.12％，Si：0.1～0.5％，Mn：1.0～2.0％，P≤0.02％，S≤0.006％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，该高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度≥300MPa，抗拉强度≥590MPa，延伸率≥20％，强塑积≥15GPa％，应变硬化指数n 90值大于0.20；更优选地，屈服强度为300～560MPa、如300～400MPa或450～560MPa，抗拉强度为590～860MPa、优选630～860MPa，延伸率为20～30％，强塑积为15～21GPa％。优选地，所述C含量为0.045～0.105％或 0.05～0.12％。优选地，所述Si含量为0.1～0.4％。优选地，所述Mn含量为1.0％～1.5％或1.2～2.0％。优选地，所述双相钢中可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.3％。优选地，所述双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，平均晶粒尺寸在1～5μm。
优选地，所述抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢通过下述工艺获得：
1)冶炼、铸造
按上述化学成分冶炼并铸造成板坯；
2)热轧、卷取
卷取温度550～680℃；
3)冷轧
冷轧压下率为40～85％；
4)快速热处理、热镀铝锌或热镀锌铝镁
冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
之后进行均热，均热温度：750～845℃，均热时间：10～60s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～200℃/s(如50～150℃/s)的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
热镀铝锌之后，以30～200℃/s(如30～150℃/s)的冷却速率快速冷却至室温，获得热镀铝锌AZ产品；或者，
热镀锌铝镁之后，以10～300℃/s(如30～180℃/s)的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
优选的，所述C含量为0.065～0.085％或0.07～0.10％。
优选的，所述Si含量为0.15～0.25％或0.1～0.4％。
优选的，所述Mn含量为1.2％～1.35％或1.5～1.8％。
优选的，所述双相钢中可含有Cr、Mo、Ti、Nb、V中的一种或两种，且 Cr+Mo+Ti+Nb+V≤0.4％，或≤0.2％。
优选的，步骤4)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s，优选29～122s。
优选的，步骤2)中，所述热轧终轧温度≥A r3。
优选的，步骤2)中，所述卷取温度为580～650℃。
优选的，步骤3)中，所述冷轧压下率为60～80％。
优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤4)中，热镀锌铝镁之后，以30～250℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
本发明所述双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，平均晶粒尺寸在1～5μm，如1～3μm。
在一些实施方案中，本发明所述抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢的化学成分质量百分比为：C：0.045～0.105％，Si：0.1～0.4％，Mn：1.0～1.5％，P≤0.02％，S≤0.006％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.3％，余量为Fe和其它不可避免的杂质。优选地，此双相钢的C含量为0.065～0.085％；优选地，此双相钢的Si含量为0.15～0.25％；优选地，此双相钢的Mn含量为1.2％～1.35％。优选地，此双相钢中可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.2％。优选地，此双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，平均晶粒尺寸在1～3μm。优选地，此双相钢的屈服强度为30～400MPa，抗拉强度为630～700MPa，延伸率为22～30％，强塑积15～20GPa％，应变硬化指数n 90值大于0.21；更优选地，此双相钢的屈服强度为304～398MPa，抗拉强度为630～698MPa，延伸率为22.3～29.4％，强塑积15.3～19.4GPa％，应变硬化指数n 90值大于0.21。优选地，此双相钢通过下述工艺获得：
a)冶炼、铸造
按上述化学成分冶炼并铸造成板坯；
b)热轧、卷取
卷取温度550～680℃；
c)冷轧
冷轧压下率为40～85％；
d)快速热处理、热镀铝锌或热镀锌铝镁
冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
之后进行均热，均热温度：750～845℃，均热时间：10～60s；均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～200℃/s(如50～150℃/s)的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
热镀铝锌之后，以30～200℃/s的冷却速率冷却至室温，获得热镀铝锌AZ产品；或者，
热镀锌铝镁之后，以30～180℃/s的冷却速率冷却至室温，获得热镀锌铝镁AM产品。
优选的，步骤d)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～122s。
优选的，步骤b)中，所述热轧终轧温度≥A r3。
优选的，步骤b)中，所述卷取温度为580～650℃。
优选的，步骤c)中，所述冷轧压下率为60～80％。
优选的，步骤d)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤d)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤d)中，所述快速加热采用两段式加热：第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃ 加热至750～845℃。
在一些实施方案中，所述抗拉强度≥590MPa的高成形性热镀铝锌活热镀锌铝镁双相钢化学成分质量百分比为：C：0.05～0.12％，Si：0.1～0.5％，Mn：1.2～2.0％，P≤0.015％，S≤0.003％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，此双相钢的C含量为0.07～0.10％。优选的，此双相钢的Si含量为0.1～0.4％。优选的，此双相钢的Mn含量为1.5～1.8％。优选的，所述双相钢中可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.4％。优选地，此双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，平均晶粒尺寸在1～5μm。优选地，此双相钢的屈服强度为470～560MPa，抗拉强度为780～860MPa，延伸率为20～25％，强塑积16～21GPa％，应变硬化指数n 90值大于0.20；更优选地，此双相钢的屈服强度为476～556MPa，抗拉强度为786～852MPa，延伸率为20.1～24.8％，强塑积16.7～20.2GPa％，应变硬化指数n 90值大于0.20。优选地，此双相钢通过下述工艺获得：
A)冶炼、铸造
按上述化学成分冶炼并铸造成板坯；
B)热轧、卷取
卷取温度550～680℃；
C)冷轧
冷轧压下率为40～85％；
D)快速热处理、热镀铝锌或热镀锌铝镁
冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
之后进行均热，均热温度：750～845℃，均热时间：10～60s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～200℃/s(如50～150℃/s)的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
热镀铝锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀铝锌AZ产品；或者，
热镀锌铝镁之后，以10～300℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
优选的，步骤D)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s。
优选的，步骤B)中，所述热轧终轧温度≥A r3。
优选的，步骤B)中，所述卷取温度为580～650℃。
优选的，步骤C)中，所述冷轧压下率为60～80％。
优选的，步骤D)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤D)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤D)中，所述快速加热采用两段式加热：第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤D)中，热镀锌铝镁之后，以30～250℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
在本发明钢的成分与工艺设计中：
C：碳是钢中最常见的强化元素，碳使钢的强度增加，塑性下降，但对成形用钢而言，需要的是低的屈服强度、高的均匀延伸率和总延伸率，故碳含量不宜过高。碳含量对钢的力学性能影响十分大，随着含碳量的升高，珠光体的数量会增加，钢的强度与硬度会大幅提高，但是其塑性与韧性会明显下降，若含碳量过高，钢中便会出现明显的网状碳化物，而网状碳化物的存在会使其强度、塑性与韧性都明显下降，钢中含碳量的升高所产生的强化效果也会显著减弱，使钢的工艺性能变差，所以在保证强度的前提下应尽量降低碳含量。
对于双相钢而言，碳元素主要影响退火过程中形成的奥氏体的体积分数，在奥氏体的形成过程当中，碳元素在奥氏体或铁素体中的扩散过程实际上起到了控制奥氏体晶粒长大的过程。随碳含量升高或临界区加热温度升高，奥氏体体积分数增加，进而冷却后所形成的马氏体相组织增加，材料的强度增加，所以综合考虑材料强韧 性匹配、快速退火过程强度的提升。本发明将含碳量限定在0.045～0.12％范围之内。
Mn：锰可以与铁形成固溶体，进而提高碳钢中铁素体与奥氏体的强度及硬度，并使钢材在热轧之后的冷却过程中获得较细小且强度较高的珠光体，而且珠光体的含量也会随着Mn含量的增加而有所增加。锰同时又是碳化物的形成元素，锰的碳化物能够溶入渗碳体，从而间接地增强珠光体的强度。锰还可以强烈增强钢的淬透性，进一步提高其强度。
对于双相钢而言，锰元素是明显影响临界区退火时奥氏体形成动力学的元素之一，锰主要影响奥氏体生成后向铁素体转变并长大的过程，以及奥氏体与铁素体的最终平衡过程。由于锰元素在奥氏体中的扩散速度远小于其在铁素体中的扩散速度，受锰扩散控制的奥氏体晶粒长大的时间较长，而锰元素在奥氏体内达到均匀分布的时间会更长。在临界区加热时，如果保温时间较短，锰元素在奥氏体内达不到均匀分布，随后冷却速率不足，就会得不到均一的马氏体岛组织。在采用快速热处理工艺生产的双相钢中(如水淬连续退火生产线)，含锰量一般较高，致使奥氏体生成后即具有较高的锰含量，保证奥氏体岛的淬透性，冷却后得到均一的马氏体岛组织和较均匀的性能。此外，锰元素扩大γ相区，降低A c1和A c3温度，因此含锰钢在同样热处理条件下将比低碳钢得到更高的马氏体体积分数。但锰含量较高时，有使钢中晶粒粗化的趋势，增加钢的过热敏感性；当熔炼浇注与热锻轧之后冷却不当时，容易使碳钢中产生白点。综合以上因素考虑，本发明将含锰量设计在1.0～2.0％范围之内。
Si：硅在铁素体或奥氏体中形成固溶体，从而增强钢的屈服强度与抗拉强度，而且硅可增大钢的冷加工变形硬化率，是合金钢中的有益元素。另外硅在硅锰钢的沿晶断口表面有着明显的富集现象，硅在晶界位置的偏聚能够减缓碳与磷沿晶界的分布，进而改善晶界的脆化状态。硅可以提高钢的强度、硬度与耐磨性，而且不会使钢的塑性明显下降。硅脱氧的能力较强，是炼钢时常用的脱氧剂，硅还能够增大钢液的流动性，所以一般钢中都含硅，但是当钢中硅的含量过高时，其塑性与韧性会显著下降。
对于双相钢而言，硅的主要影响是降低给定退火时间条件下最终平衡时的奥氏体体积分数。硅对奥氏体长大速率没有明显影响，但对奥氏体的形成形态和分布有明显影响。因此，本发明将含硅量确定在0.1～0.5％范围之内。
Cr：铬在钢中的主要作用是提高淬透性。使钢经淬火回火后具有较好的综合力 学性能。铬与铁形成连续固溶体，缩小奥氏体相区城。铬与碳形成多种碳化物，与碳的亲和力大于铁和锰元素。铬与铁可形成金属间化合物σ相(FeCr)，铬使珠光体中碳的浓度及奥氏体中碳的极限溶解度减少；铬减缓奥氏体的分解速度，显著提高钢的淬透性。但亦增加钢的回火脆性倾向。铬元素可提高钢的强度和硬度同时加入其他合金元素时，效果较显著。由于Cr提高了钢在空冷时的淬火能力，因而对钢的焊接性能有不利的影响。但是在含铬量小于0.3％时，对焊接性的不利影响可以忽略；大于此含量时，容易在焊接时产生裂纹和夹渣等缺陷。当Cr与其他合金元素同时存在(如和V共存)时，Cr对焊接性的不利影响大大减小。如当Cr、Mo、V等元素同时存在于钢中时，即使含Cr量达到1.7％，对钢的焊接性能尚无显著的不利影响。本发明中铬元素为有益且非必要添加元素，考虑成本增加等因素添加量不宜过多。在一些实施方案中，Cr含量≤0.35％。
Mo：钼元素能抑制铁的自扩散和其他元素的扩散速度。Mo原子半径比α-Fe原子大，当Mo溶解在α固溶体时，使固溶体发生强烈的晶格畸变，同时Mo能增加晶格原子键引力，提高α铁素体的再结晶温度。Mo在珠光体型、铁素体型、马氏体型钢中，甚至在高合金奥氏体钢中的强化作用也十分明显。Mo在钢中的良好作用还需视与钢中其他合金元素间的相互作用而定。在钢中加入强碳化物形成元素V、Nb、Ti时，Mo的固溶强化作用更加显著。这是因为当强碳化物形成元素与C结合成稳定的碳化物时，能促进Mo更有效地溶入固溶体中，从而更有利于钢的热强性提高。加入Mo还可以增加钢的淬透性，但效果没有C和Cr显著。Mo会抑制珠光体区的转变，使中温区转变加快，因而含Mo钢在冷却速度较大的情况下也能形成一定数量的贝氏体，并且消除铁素体的形成，这是Mo对低合金耐热钢热强性产生有利影响的原因之一。Mo还能显著降低钢的热脆倾向，并减小珠光体球化速度。当Mo含量在0.15％以下时，对钢的焊接性能无不利的影响。本发明中钼元素为有益且非必要添加元素，考虑成本增加等因素添加量不宜过多。在一些实施方案中，Mo含量≤0.2％
微合金元素Ti、Nb、V：在钢中添加微量的微合金元素Nb、V、Ti，可保证钢在碳当量较低的情况下，通过其碳、氮化物质点(尺寸小于5nm)的弥散析出及Nb、V、Ti的固溶，细化晶粒，极大地提高钢的强度、韧性，特别是低温韧性，使钢具有良好的可焊性、使用性。Nb、V、Ti是碳化物和氮化物的形成元素，这些元素在比较低的浓度下就能满足这种要求Nb、V、Ti为强碳化物形成元素，常温时，在钢 中大部分以碳化物、氮化物、碳氮化物形式存在，少部分固溶在铁素体中。加入Nb、V、Ti可以阻止奥氏体晶粒长大，提高钢的粗化温度。这是由于它们的碳、氮化物弥散的小颗粒能对奥氏体晶界起固定作用，阻碍奥氏体晶界的迁移，提高奥氏体再结晶温度，可扩大未再结晶区，亦即阻止了奥氏体晶粒长大。在钢中添加微量的Nb、V、Ti，一方面，可在减少碳当量含量的同时提高强度，从而提高钢的焊接性能；另一方面，将不纯物质如氧、氮、硫等固定起来，从而改善钢的可焊性；其次，由于其微观质点的作用，例如TiN在高温下的未溶解性，可阻止热影响区晶粒的粗化，提高热影响区的韧性，从而改善钢的焊接性能。本发明中微合金元素为有益且非必要添加元素，考虑成本增加等因素添加量不宜过多。在一些实施方案中，Ti含量≤0.04％。在一些实施方案中，Nb含量≤0.05％。在一些实施方案中，V含量≤0.05％。
本发明通过快速加热、短时保温和快速冷却的快速热处理工艺控制连续热处理工艺中变形组织的回复、再结晶、奥氏体相变及晶粒长大等过程，在冷却过程中不仅形成铁素体基体相，且产生各种强化相和相内的成分梯度分布，最终获得细小的铁素体组织及多形态的强化相组织，使材料获得较佳的强韧性配合，降低合金成本和各工序制造难度，提高相同强度级别钢种的焊接性能等使用性能。
具体原理在于：加热过程不同温度阶段采用不同加热速率，低温段主要发生变形组织的回复，可采用相对低的加热速率以降低能耗；高温段主要发生不同相组织的再结晶和晶粒长大，必须要采用相对高的加热速率来缩短组织在高温区间的停留时间才能确保晶粒无法长大。通过控制加热过程中的加热速率抑制加热过程中变形组织的回复及铁素体再结晶过程，使再结晶过程与奥氏体相变过程重叠，增加了再结晶晶粒和奥氏体晶粒的形核点，最终细化晶粒。通过短时保温和快速冷却，缩短均热过程晶粒长大的时间，确保晶粒组织细小、均匀分布。
本发明所述的抗拉强度≥590MPa的高成形性热镀铝锌及热镀锌铝镁双相钢的快速热处理热镀制造方法包括以下步骤：
1)冶炼、铸造
按上述化学成分冶炼并铸造成板坯；
2)热轧、卷取
卷取温度550～680℃；
3)冷轧
冷轧压下率40～85％，冷轧后获得轧硬态带钢或钢板；
4)快速热处理、热镀铝锌或热镀锌铝镁
A)快速加热
将冷轧带钢或钢板由室温快速加热至750～845℃奥氏体和铁素体两相区目标温度，所述快速加热采用一段式或两段式；
采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
B)均热
在奥氏体和铁素体两相区目标温度750～845℃进行均热，均热时间为10～60s；
C)冷却
带钢或钢板均热结束后以5～15℃/s冷却速率缓冷至670～770℃；随后以50～200℃/s(如50～150℃/s)冷却速率快速冷却至580～600℃；
D)热镀铝锌或热镀锌铝镁
将带钢或钢板快速冷却至580～600℃后浸入锌锅进行热镀铝锌或热镀锌铝镁；
E)热镀铝锌之后，以30～200℃/s(如30～150℃/s)的冷却速率快速冷却至室温，
获得热镀铝锌AZ产品；或者，
热镀锌铝镁之后，以10～300℃/s(30～180℃/s)的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
优选的，所述快速热处理热镀铝锌或热镀锌铝镁全过程用时为29～159s，如29～122s。
优选地，在一些实施方案中，所述方法用于制备化学成分质量百分比如下所述的前文任一实施方案所述的双相钢：C：0.045～0.105％，Si：0.1～0.4％，Mn：1.0～1.5％，P≤0.02％，S≤0.006％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.3％，余量为Fe和其它不可避免的杂质；其中，热镀铝锌之后，以30～200℃/s的冷却速率冷却至室温，获得热镀铝锌AZ产品，或者，热镀锌铝镁之后，以30～180℃/s的冷却速率冷却至室温，获得热镀锌铝镁AM产品。优选地，此方法所述快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～122s。
优选地，在一些实施方案中，所述方法用于制备化学成分质量百分比如下所述的前文任一项实施方案所述的双相钢：C：0.05～0.12％，Si：0.1～0.5％，Mn：1.2～2.0％，P≤0.015％，S≤0.003％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；其中，热镀铝锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀铝锌AZ产品，或者，热镀锌铝镁之后，以10～300℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。优选地，此方法所述快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s。
优选的，步骤2)中，所述热轧终轧温度≥A r3。
优选的，步骤2)中，所述卷取温度为580～650℃。
优选的，步骤3)中，所述冷轧压下率为60～80％。
优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤4)中，所述快速加热采用两段式加热，第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤4)中，所述快速加热采用两段式加热，第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃。
优选的，步骤4)中，所述快速加热最终温度为770～830℃。
优选的，步骤4)均热过程中，带钢或钢板加热至所述奥氏体和铁素体两相区目标温度后，保持温度不变进行均热。
优选的，步骤4)均热过程中，带钢或钢板在均热时间段内进行小幅度升温或小幅度降温，升温后温度不超过845℃，降温后温度不低于750℃。
优选的，所述均热时间为10～40s。
在本发明所述的590MPa级别高成形性热镀铝锌或热镀锌铝镁双相钢的快速热处理热镀制造方法中：
1、加热速度控制
连续加热过程的再结晶动力学可以由受加热速率影响的关系式来定量描述，连续加热过程中铁素体再结晶体积分数与温度T的函数关系式为：
Figure PCTCN2022084543-appb-000001
其中，X(t)为铁素体再结晶体积分数；n为Avrami指数，与相变机制有关，取决于再结晶形核率的衰减周期，一般在1～4的范围内取值；T为热处理温度；T star为再结晶开始温度；β是加热速率；b(T)由下式所获得：
b＝b 0exp(-Q/RT)
从以上公式及有关实验数据可以得出，随加热速率增加，再结晶开始温度(T star)及结束温度(T fin)均升高；加热速率在50℃/s以上时，奥氏体相变与再结晶过程将重叠，再结晶温度升高至两相区温度，加热速率越快，铁素体再结晶温度也越高。
传统热处理过程采用慢速加热，该条件下变形基体依次发生回复、再结晶及晶粒长大，而后发生铁素体向奥氏体的相转变，相变形核点主要集中在已经长大的铁素体晶界处，形核率较低，最终得到的晶粒组织比较粗大。
快速加热条件下，变形基体还没有完成回复就开始发生铁素体向奥氏体的相转变，或再结晶刚刚完成，晶粒还没有长大就发生奥氏体相变，由于刚刚完成再结晶时晶粒细小、晶界面积大，因此形核率显著提高，奥氏体晶粒明显细化。特别是铁素体再结晶与奥氏体相变过程发生重叠后，由于铁素体晶体内保留了大量位错等晶体缺陷，为奥氏体提供了大量的形核点，使得奥氏体呈现爆发式形核，奥氏体晶粒进一步细化。同时保留下来的高密度位错线缺陷也成为碳原子高速扩散的通道，使得每一个奥氏体晶粒都能快速生成并长大，因此增大奥氏体体积分数。
快速加热过程中精细控制组织演变、合金元素和各相组分分布，为后续均热过程奥氏体组织长大，以及各合金成分分布及快速冷却过程奥氏体向马氏体相转变奠定了良好的基础。最终才能获得具有细化晶粒、合理的元素及各相分布的最终产品组织。综合考虑快速加热细化晶粒的效果、制造成本以及可制造性等因素，本发明将一段式快速加热时加热速率定为15～500℃/s，采用两段式快速加热时加热速率定为30～500℃/s。。
由于不同温度区间范围内，快速加热对材料的回复、再结晶和晶粒长大等组织演变过程所产生的影响不同，为获得最优的组织控制，因此不同的加热温度区间其优选的加热速率也不相同：从20℃到550～650℃，加热速率对回复过程的影响最大，控制加热速率为15～500℃/s，进一步优选为30～500℃/s；加热温度从550～650℃到奥氏体化温度750～845℃，加热速率对晶粒长大过程影响最大，控制加热速率为 50～500℃/s；进一步优选为80～500℃/s。
2、均热温度控制
均热温度通常取决于C含量，本发明双相钢中C含量为0.045～0.12％，本发明钢的A C1和A C3分别在730℃和870℃左右。本发明的快速热处理工艺中将带钢加热到A C1到A C3之间进行均热，利用快速加热技术在未充分再结晶的铁素体中保留大量的位错，为奥氏体转变提供了更大的形核驱动力，所以较传统连续退火工艺，本发明的快速热处理方法可获得更多更细小的奥氏体组织。
本发明对于均热温度的控制，率先提出均热温度在一定范围内进行升高和降低：即均热带温倾斜升温和倾斜降温，但均热温度必须保持在一定范围之内。这样做的好处在于：在两相区温度范围内快速升降温过程，实际上就是进一步增加过热度和过冷度，便于快速相转变过程，当升降温幅度足够大、升降温速率也足够大时，可以通过反复的铁素体向奥氏体相转变以及奥氏体向铁素体相转变进一步细化晶粒，同时对碳化物形成及合金元素的均匀分布也起到一定的影响，最终形成更细小的组织及具有均匀分布的合金元素。
冷轧后双相钢中有大量未溶解的细小均匀分布的碳化物，在加热过程中，能够对奥氏体晶粒的长大起到机械阻碍的作用，有利于细化高强钢的晶粒度。但是如果均热温度过高，就会使未溶解的碳化物数目大量减少，削弱这种阻碍作用，增强晶粒的长大倾向，进而降低钢的强度。当未溶碳化物的数量过多时，又有可能引起聚集，造成局部化学成分的不均匀分布，该聚集处的含碳量过高时，还会引发局部过热。所以理想情况下，钢中应该均匀分布着少量细小的颗粒状未溶碳化物，这样既可以防止奥氏体晶粒异常长大，又能够相应地提高基体中的各合金元素的含量，达到改善合金钢的强度与韧性等力学性能的目的。
均热温度的选取还应以获得细小均匀的奥氏体晶粒为目的，避免奥氏体晶粒粗大，以达到在冷却之后能够得到细小的马氏体组织的目的。过高的均热温度会使奥氏体晶粒粗大，快冷后获得的马氏体组织也会较粗大，使钢的力学性能不佳；还会增加残余奥氏体的数量、减少马氏体的数量，降低钢的硬度与耐磨性。过低的均热温度，又会使奥氏体溶入的碳以及合金元素不足，令奥氏体中合金元素浓度分布不均，使钢的淬透性大幅降低，对钢的力学性能造成不利影响。亚共析钢的均热温度应该为Ac3+30～50℃。对于超高强度钢来说，存在碳化物形成元素，会阻碍碳化物的转变，所以均热温度可以适当的提高。综合以上因素，本发明选取750～845℃作 为均热温度，以期获得更理想更合理的最终组织。
3、均热时间控制
均热时间的影响因素也取决于钢中碳以及合金元素的含量，当其含量升高时，不仅会导致钢的导热性降低，而且因为合金元素比碳元素的扩散速度更慢，合金元素会明显延滞钢的组织转变，这时就要适当延长保温时间。由于本工艺采用快速加热，在两相区因材料含有大量残余位错，因此为奥氏体形成提供大量的形核点，且为碳原子提供了快速扩散通道，所以奥氏体可以极快的形成，而且均热保温时间越短碳原子扩散距离越短，奥氏体内碳浓度梯度越大，因此最后保留下来的残余奥氏体碳含量越多；但是如果均热保温时间过短，会使钢中合金元素分布不均而且会导致奥氏体化不充分；保温时间过长又容易导致奥氏体晶粒粗大。均热保温时间的影响因素也取决于钢中碳以及合金元素的含量，当其含量升高时，不仅会导致钢的导热性降低，而且因为合金元素比碳元素的扩散速度更慢，合金元素会明显延滞钢的组织转变，这时就要适当延长保温时间。所以均热时间的控制需严格结合均热温度、快速冷却及快速加热过程综合考虑制定，才能最终获得理想的组织和元素分布。综上，本发明将保温时间定为10～60s。
4、快速冷却速度控制
快速冷却过程控制需结合前期加热和均热过程中各组织演变结果及合金元扩散分布结果等综合因素，确保最终获得理想的各相组织及元素合理分布的材料组织。
为了获得足够的马氏体强化相，淬火时试样的冷速必须大于临界冷却速度才能够得到马氏体组织，而临界冷却速度主要取决于材料成分，本发明中的Si含量为0.1～0.5％，Mn含量为1.0～2.0％，含量相对较高，所以Si和Mn很大程度加强了双相钢的淬透性，从而降低了临界冷却速度。同时冷却速率还需综合考虑加热过程和均热过程的组织演变及合金扩散分布结果，以最终获得合理的各相组织分布及合金元素分布。冷却速率太低无法获得马氏体组织，结果导致强度下降，力学性能无法满足要求；但是太大的冷速又会产生较大的淬火应力(即组织应力与热应力)，容易导致试样变形甚至开裂。所以本发明将快速冷却速度设置为50～150℃/s。
对于高强度的热镀铝锌及热镀锌铝镁产品而言，快速热处理工艺由于减少了带钢在高温炉内的停留时间，因此在热处理过程中合金元素在高强度带钢表面的富集量显著减少，有利于高强度热镀锌产品可镀性的改善，因而有利于高强度热镀锌产 品表面漏镀缺陷的减少和耐蚀性能的提高，从而能提高成材率；另外由于产品晶粒的细化和材料合金含量的减少使得采用本发明技术得到的双相钢产品的扩孔性能和弯折性能等加工成形性能及焊接性能等用户使用性能也有所提高。
快速热处理工艺技术使得加热过程和均热过程时间减少，炉子长度缩短(较传统连续退火炉至少能缩短三分之一)，炉辊数量显著减少，使得炉内产生表面缺陷的几率减少，因此产品表面质量将显著提高。
本发明相对于传统技术所具有的优点：
(1)本发明通过快速热处理抑制热处理过程中变形组织的回复及铁素体再结晶过程，使再结晶过程与奥氏体相变过程重叠，增加再结晶晶粒和奥氏体晶粒的形核点，缩短晶粒长大时间，所获得的双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，快速热处理后组织中的细小马氏体，其特征是具有块状，条状，颗粒状等多种形态，且分布更加均匀，从而双相钢产品可获得良好的强塑性匹配。
(2)相比于传统连续退火热镀锌方式所得的热镀锌双相钢，在前工序制造条件不变的前提下，通过本发明快速热处理后得到的双相钢的平均晶粒尺寸为1～5μm，可获得良好的细晶强化的效果。其屈服强度为300～560MPa，抗拉强度为590～860MPa，延伸率为20～30％，强塑积为15～21GPa％，应变硬化指数n 90值大于0.20。
(3)根据本发明所述的低碳低合金高成形性低碳低合金热镀锌双相钢快速热处理工艺，热处理全过程用时可缩短至29～159s，大大降低了整个热处理工艺过程的时间(传统连续退火工艺时间通常在300～480s)，显著提高了生产效率、减少了能耗，降低了生产成本。
(4)在生产成本和制造难度方面，相比于传统的双相钢及其热处理工艺，本发明的快速热处理方法缩短了连续热镀锌退火炉加热段和均热段的长度(较传统连续退火炉至少能缩短三分之一)、时间及整个热处理工序时间，可节能减排降耗，并显著降低炉子设备一次性投资，显著降低生产运行成本和设备维护成本；另外通过快速热处理生产相同强度等级的产品可以降低合金含量，降低热处理及前工序的生产成本，降低热处理之前各工序的制造难度。
(5)在产品质量方面，相比于传统连续退火处理得到的双相钢，由于快速热处理工艺技术使得加热过程和均热过程时间减少炉子长度缩短，炉辊数量显著减少，使得炉内产生表面缺陷的几率减少，因此产品表面质量将显著提高；对于高强 度的热镀铝锌及热镀锌铝镁产品而言，快速热处理工艺由于减少了带钢在高温炉内的停留时间，因此在热处理过程中合金元素在高强度带钢表面的富集量显著减少，有利于产品可镀性的改善，因而有利于高强度热镀铝锌及热镀锌铝镁产品表面漏镀缺陷的减少和耐蚀性能的提高，从而能提高成材率；另外由于产品晶粒的细化和材料合金含量的减少使得采用本发明技术得到的双相钢产品的扩孔性能和弯折性能等加工成形性能及焊接性能等用户使用性能也有所提高。
(6)相比于传统热处理得到的热镀双相钢，本发明技术得到的热镀热镀铝锌及热镀锌铝镁双相钢的成形、焊接、涂装、耐蚀性能等用户使用性能也有所提高，其中耐蚀性是传统GI双相钢的3～8倍。
综上所述，通过本发明得到的高成形性热镀锌双相钢对新一代轻量化汽车、火车、船舶、飞机等交通运输工具的发展和相应工业的健康发展以及先进制造业的健康发展均具有重要价值。
附图说明
图1是本发明实施例一试验钢A按实施例1(一段式加热)所生产的热镀锌铝镁双相钢(AM)显微组织图片。
图2是本发明实施例一试验钢A按实施例1(两段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片。
图3是本发明实施例一试验钢A按传统工艺1所生产的热镀热镀铝锌双相钢(AZ)显微组织图片。
图4是本发明实施例一试验钢I按实施例3(两段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片。
图5是本发明实施例一试验钢C按实施例15(两段式加热)所生产的热镀锌铝镁双相钢(AM)显微组织图片。
图6是本发明实施例二试验钢D按实施例4(一段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片。
图7是本发明实施例二试验钢D按传统工艺4所生产的热镀铝锌双相钢(AZ)显微组织图片。
图8是本发明实施例二试验钢N按实施例15(两段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片。
图9是本发明实施例二试验钢E按实施例17(两段式加热)所生产的热镀锌铝镁双相钢(AM)显微组织图片。
具体实施方式
下面结合实施例和附图对本发明作进一步说明，本实施例以本发明技术方案为前提进行实施，给出了详细的实施方式和具体操作过程，但本发明的保护范围不限于下述的实施例。
实施例中，屈服强度、抗拉强度和延伸率依据《GB/T228.1-2010金属材料拉伸试验第1部分：室温试验方法》进行，采用P7号试样沿横向进行测试。n 90依据《GB/T228.1-2010金属材料拉伸试验第1部分：室温试验方法》进行，采用P7号试样沿横向进行测试，依据《GBT 5028-2008金属材料薄板和薄带拉伸应变硬化指数(n值)的测定方法》获得n 90值。
实施例一
本实施例试验钢的成分参见表1，本实施例按照一段式快速热处理实施例及传统工艺的具体参数参见表2，本实施例按照二段式快速热处理实施例及传统工艺的具体参数参见表3，表4为本实施例试验钢成分按一段式加热实施例及传统工艺所得双相钢的主要性能，表5为本实施例试验钢成分按二段式加热实施例及传统工艺制备所得双相钢的主要性能。
从表1～表5可以看出，通过本发明的方法，可降低同级别钢中的合金含量，细化晶粒、获得材料组织构成及强度和韧性的匹配。通过本发明的方法获得的双相钢屈服强度为304～398MPa，抗拉强度为630～698MPa，延伸率为22.3～29.4％，强塑积15.3-19.4GPa％，应变硬化指数n 90值大于0.21。
图1是本实施例试验钢A按本发明实施例1(一段式加热)所生产的热镀锌铝镁双相钢(AM)显微组织图片，图2、图3为典型成分A钢经过实施例1和对比传统工艺例1的组织图(两段式加热)，图4是本实施例I钢通过实施例3(两段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片，图5是本实施例C钢通过实施例15中(两段式加热)传统加热速率下所生产的热镀锌铝镁双相钢(AM)显微组织图片。
从图1～图5可以看出，全部材料组织均由铁素体、马氏体及少量碳化物组成。 如图3所示，采用传统工艺处理的组织特点是：晶粒粗大，且存在一定的带状组织，马氏体及碳化物沿铁素体晶界呈网状分布，铁素体晶粒相对粗大，铁素体及马氏体两相组织分布不均匀。
Figure PCTCN2022084543-appb-000002
Figure PCTCN2022084543-appb-000003
Figure PCTCN2022084543-appb-000004
Figure PCTCN2022084543-appb-000005
表4
Figure PCTCN2022084543-appb-000006
Figure PCTCN2022084543-appb-000007
表5
Figure PCTCN2022084543-appb-000008
Figure PCTCN2022084543-appb-000009
Figure PCTCN2022084543-appb-000010
实施例二
本实施例试验钢的成分参见表6，本实施例按照一段式快速热处理实施例及传统工艺的具体参数参见表7，本实施例按照二段式快速热处理实施例及传统工艺的具体参数参见表8，表9为本实施例试验钢成分按一段式加热所得实施例及传统工艺所得双相钢的主要性能，表10为本实施例试验钢成分按二段式加热所得实施例及传统工艺制备所得双相钢的主要性能。
从表6～表10可以看出，通过本发明的方法，可降低同级别钢中的合金含量，细化晶粒、获得材料组织构成及强度和韧性的匹配。通过本发明的方法获得的双相钢屈服强度为476～556MPa，抗拉强度为786～852MPa，延伸率为20.1～24.8％，强塑积16.7～20.2GPa％，应变硬化指数n 90值大于0.20。
图6、图7为本实施例实验钢D经过实施例4和对比传统工艺例4(一段式加热)的热镀铝锌双相钢(AZ)显微组织图，图8是本实施例实验钢N通过实施例15(两段式加热)所生产的热镀铝锌双相钢(AZ)显微组织图片，图9是本实施例实验钢E通过实施例17(两段式加热)所生产的热镀锌铝镁双相钢(AM)显微组织图片。
Figure PCTCN2022084543-appb-000011
Figure PCTCN2022084543-appb-000012
Figure PCTCN2022084543-appb-000013
Figure PCTCN2022084543-appb-000014
表9
Figure PCTCN2022084543-appb-000015
表10
Figure PCTCN2022084543-appb-000016
通过本发明工艺处理获得的双相钢组织特点：铁素体、马氏体晶粒组织及碳化物都非常细小且均匀分布于基体中，这对提高材料强度和塑性都是非常有利的。因此本发明的双相钢制备方法可细化晶粒，使材料各相组织均匀分布于基体中，进而改善材料组织，提高材料性能。
本发明通过采用快速加热和快速冷却工艺对传统连续退火机组进行工艺改造，使其实现快速热处理工艺，可以极大的缩短传统连续退火炉加热段及均热段的长度，提高传统连续退火机组的生产效率，降低生产成本及能耗，减少连续退火炉的炉辊数量，这可以提高带钢表面质量控制能力，获得高表面质量的带钢产品；同时通过建立采用快速热处理工艺技术的新型连续退火机组，可实现机组短小精悍、产品规格品种过渡灵活、调控能力强等优点；对材料而言则可细化带钢晶粒，进一步提高材料强度，降低合金成本及热处理前工序的制造难度，提高材料的成形、焊接等用户使用性能。
综上所述，本发明通过采用快速热处理工艺，对冷轧带钢的连续退火工艺技术进步产生了极大的促进作用，冷轧带钢从室温开始到最后完成奥氏体化过程可望在十几秒甚至几秒内完成，大大缩短了连续退火炉子加热段长度，便于提高连续退火机组的速度和生产效率，显著减少连续退火机组炉内辊子数目，对于机组速度在180米/分左右的快速热处理产线其高温炉段内的辊子数目不超过10根，可明显提高带钢表面质量。同时，在极短时间内所完成的再结晶和奥氏体化过程的快速热处理工艺方法也将提供更加灵活及柔性化的高强钢组织设计方法，进而在无需改变合金成分以及轧制工艺等前工序条件的前提下改善材料组织，提高材料性能。
以双相钢为代表的高耐腐性镀层先进高强钢有着广阔的应用前景，而快速热处理技术又有着巨大的开发价值，两者的结合必将会为双相钢的开发和生产提供更大的空间。

Claims (15)

  1. 抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢，其化学成分质量百分比为：C：0.045～0.12％，Si：0.1～0.5％，Mn：1.0～2.0％，P≤0.02％，S≤0.006％，Al：0.02～0.055％，任选含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度≥300MPa，抗拉强度≥590MPa，延伸率≥20％，强塑积≥15GPa％，应变硬化指数n 90值大于0.20；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的金相组织为均匀分布的铁素体和马氏体双相组织，平均晶粒尺寸为1～5μm；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢通过下述工艺获得：
    1)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    2)热轧、卷取
    卷取温度550～680℃；
    3)冷轧
    冷轧压下率为40～85％；
    4)快速热处理、热镀铝锌或热镀锌铝镁
    冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
    采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
    采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
    之后进行均热，均热温度：750～845℃，均热时间：10～60s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～200℃/s(如50～150℃/s)的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
    热镀铝锌之后，以30～200℃/s(如30～150℃/s)的冷却速率快速冷却至室温，获得热镀铝锌AZ产品；或者，
    热镀锌铝镁之后，以10～300℃/s(如30～180℃/s)的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
  2. 如权利要求1所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，
    所述C含量为0.045～0.105％或0.05～0.12％，优选为0.065～0.085％或0.07～0.10％；和/或
    所述Si含量为0.15～0.25％或0.1～0.4％；和/或
    所述Mn含量为1.0％～1.5％或1.2～2.0％，优选为1.2％～1.35％或1.5～1.8％；和/或
    所述Cr+Mo+Ti+Nb+V≤0.4％，或≤0.2％。
  3. 如权利要求1或2所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度为300～560MPa、如300～400MPa或450～560MPa，抗拉强度为590～860MPa、优选630～860MPa，延伸率为20～30％，强塑积为15～21GPa％。
  4. 如权利要求1-3中任一项所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，
    步骤2)中，所述热轧终轧温度≥A r3；和/或
    步骤2)中，所述卷取温度为580～650℃；和/或
    步骤3)中，所述冷轧压下率为60～80％。
  5. 如权利要求1-4中任一项所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，
    步骤4)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s，优选29～122s；和/或
    步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    步骤4)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃；优选地，第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃；和/或
    步骤4)中，热镀锌铝镁之后，以30～250℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
  6. 如权利要求1所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于， 所述抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢的化学成分质量百分比为：C：0.045～0.105％，Si：0.1～0.4％，Mn：1.0～1.5％，P≤0.02％，S≤0.006％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.3％，余量为Fe和其它不可避免的杂质；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的C含量为0.065～0.085％；优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的Si含量为0.15～0.25％；优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的Mn含量为1.2％～1.35％；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢中，Cr+Mo+Ti+Nb+V≤0.2％；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的平均晶粒尺寸为1～3μm。
  7. 如权利要求6所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度为30～400MPa，抗拉强度为630～700MPa，延伸率为22～30％，强塑积15～20GPa％，应变硬化指数n 90值大于0.21；优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度为304～398MPa，抗拉强度为630～698MPa，延伸率为22.3～29.4％，强塑积15.3～19.4GPa％，应变硬化指数n 90值大于0.21。
  8. 如权利要求6或7所述的所述高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，所述双相钢通过下述工艺获得：
    a)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    b)热轧、卷取
    卷取温度550～680℃；
    c)冷轧
    冷轧压下率为40～85％；
    d)快速热处理、热镀铝锌或热镀锌铝镁
    冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
    采用一段式快速加热时，加热速率为50～500℃/s；
    采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至 550～650℃，第二段以50～500℃/s的加热速率从550～650℃加热至750～845℃；
    之后进行均热，均热温度：750～845℃，均热时间：10～60s；均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～150℃/s的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
    热镀铝锌之后，以30～200℃/s的冷却速率冷却至室温，获得热镀铝锌AZ产品；或者，
    热镀锌铝镁之后，以30～180℃/s的冷却速率冷却至室温，获得热镀锌铝镁AM产品。
  9. 如权利要求8所述的所述高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，
    步骤d)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～122s；和/或
    步骤b)中，所述热轧终轧温度≥A r3；和/或
    步骤b)中，所述卷取温度为580～650℃；和/或
    步骤c)中，所述冷轧压下率为60～80％；和/或
    步骤d)中，所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    步骤d)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃；优选地，第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃。
  10. 如权利要求1所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，所述抗拉强度≥590MPa的高成形性热镀铝锌或热镀锌铝镁双相钢的化学成分质量百分比为：C：0.05～0.12％，Si：0.1～0.5％，Mn：1.2～2.0％，P≤0.015％，S≤0.003％，Al：0.02～0.055％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的C含量为0.07～0.10％；优选的，所述高成形性热镀铝锌或热镀锌铝镁双相钢的Si含量为0.1～0.4％；优选的，所述高成形性热镀铝锌或热镀锌铝镁双相钢的Mn含量为1.5～1.8％；
    优选的，Cr+Mo+Ti+Nb+V≤0.4％；
    优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度为 470～560MPa，抗拉强度为780～860MPa，延伸率为20～25％，强塑积16～21GPa％，应变硬化指数n 90值大于0.20；更优选地，所述高成形性热镀铝锌或热镀锌铝镁双相钢的屈服强度为476～556MPa，抗拉强度为786～852MPa，延伸率为20.1～24.8％，强塑积16.7～20.2GPa％，应变硬化指数n 90值大于0.20。
  11. 如权利要求10所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，所述高成形性热镀铝锌或热镀锌铝镁双相钢通过下述工艺获得：
    A)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    B)热轧、卷取
    卷取温度550～680℃；
    C)冷轧
    冷轧压下率为40～85％；
    D)快速热处理、热镀铝锌或热镀锌铝镁
    冷轧后的钢板快速加热至750～845℃，所述快速加热采用一段式或两段式；
    采用一段式快速加热时，加热速率为50～500℃/s；
    采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以50～500℃/s的加热速率从550～650℃加热至750～845℃；
    之后进行均热，均热温度：750～845℃，均热时间：10～60s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以50～150℃/s的冷却速率快速冷却至580～600℃，浸入锌锅进行热镀铝锌或热镀锌铝镁；
    热镀铝锌之后，以30～150℃/s的冷却速率冷却至室温，获得热镀铝锌AZ产品；或者，
    热镀锌铝镁之后，以10～300℃/s的冷却速率冷却至室温，获得热镀锌铝镁AM产品。
    优选的，步骤D)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s。
  12. 如权利要求11所述的高成形性热镀铝锌或热镀锌铝镁双相钢，其特征在于，
    步骤B)中，所述热轧终轧温度≥A r3；和/或
    步骤B)中，所述卷取温度为580～650℃；和/或
    步骤C)中，所述冷轧压下率为60～80％；和/或
    步骤D)中，所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    步骤D)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃；优选地，第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃；和/或
    步骤D)中，热镀锌铝镁之后，以30～250℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
  13. 权利要求1所述的抗拉强度≥590MPa的高成形性热镀铝锌及热镀锌铝镁双相钢的快速热处理热镀制造方法，其特征在于，所述方法包括以下步骤：
    1)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    2)热轧、卷取
    卷取温度550～680℃；
    3)冷轧
    冷轧压下率40～85％，冷轧后获得轧硬态带钢或钢板；
    4)快速热处理、热镀铝锌或热镀锌铝镁
    a)快速加热
    将冷轧带钢或钢板由室温快速加热至750～845℃奥氏体和铁素体两相区目标温度，所述快速加热采用一段式或两段式；
    采用一段式快速加热时，加热速率为15～500℃/s(如50～500℃/s)；
    采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～650℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～650℃加热至750～845℃；
    b)均热
    在奥氏体和铁素体两相区目标温度750～845℃进行均热，均热时间为10～60s；
    c)冷却
    带钢或钢板均热结束后以5～15℃/s冷却速率缓冷至670～770℃；随后以50～200℃/s(如50～150℃/s)冷却速率快速冷却至580～600℃；
    d)热镀铝锌或热镀锌铝镁
    将带钢或钢板快速冷却至580～600℃后浸入锌锅进行热镀铝锌或热镀锌铝镁；
    e)热镀铝锌之后，以30～200℃/s(如30～150℃/s)的冷却速率快速冷却至室温，获得热镀铝锌AZ产品；或者，
    热镀锌铝镁之后，以10～300℃/s(30～180℃/s)的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品。
  14. 如权利要求13所述的方法，其特征在于，
    步骤4)所述的快速热处理和热镀铝锌或热镀锌铝镁全过程用时为29～159s，如29～122s；和/或
    步骤2)中，所述热轧终轧温度≥A r3；和/或
    步骤2)中，所述卷取温度为580～650℃；和/或
    步骤3)中，所述冷轧压下率为60～80％；和/或
    步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    步骤4)中，所述快速加热采用两段式加热，第一段以15～300℃/s的加热速率从室温加热至550～650℃，第二段以50～300℃/s的加热速率从550～650℃加热至750～845℃；优选地，第一段以30～300℃/s的加热速率从室温加热至550～650℃，第二段以80～300℃/s的加热速率从550～650℃加热至750～845℃；和/或
    步骤4)中，所述快速加热最终温度为770～830℃；和/或
    步骤4)均热过程中，带钢或钢板加热至所述奥氏体和铁素体两相区目标温度后，保持温度不变进行均热；和/或
    步骤4)均热过程中，带钢或钢板在均热时间段内进行小幅度升温或小幅度降温，升温后温度不超过845℃，降温后温度不低于750℃；和/或
    所述均热时间为10～40s。
  15. 如权利要求13或14所述的方法，其特征在于，
    所述方法用于制备权利要求6所述的双相钢；其中，热镀铝锌之后，以30～200℃/s的冷却速率冷却至室温，获得热镀铝锌AZ产品，或者，热镀锌铝镁之后，以30～180℃/s的冷却速率冷却至室温，获得热镀锌铝镁AM产品；优选地，此方法所述快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～122s；或
    所述方法用于制备权利要求10所述的双相钢；其中，热镀铝锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀铝锌AZ产品，或者，热镀锌铝镁之后，以10～300℃/s的冷却速率快速冷却至室温，获得热镀锌铝镁AM产品；优 选地，此方法所述快速热处理和热镀铝锌或热镀锌铝镁全过程用时29～159s。
PCT/CN2022/084543 2021-04-02 2022-03-31 高成形性热镀铝锌或热镀锌铝镁双相钢及其快速热处理热镀制造方法 WO2022206917A1 (zh)

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