WO2022206912A1 - 抗拉强度≥980MPa的低碳低合金TRIP钢或热镀锌TRIP钢及其制造方法 - Google Patents

抗拉强度≥980MPa的低碳低合金TRIP钢或热镀锌TRIP钢及其制造方法 Download PDF

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WO2022206912A1
WO2022206912A1 PCT/CN2022/084524 CN2022084524W WO2022206912A1 WO 2022206912 A1 WO2022206912 A1 WO 2022206912A1 CN 2022084524 W CN2022084524 W CN 2022084524W WO 2022206912 A1 WO2022206912 A1 WO 2022206912A1
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low
steel
hot
alloy
trip steel
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PCT/CN2022/084524
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French (fr)
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李俊
王健
张利祥
张宝平
张理扬
陈培林
陈云鹏
孟庆格
杜小峰
马新建
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宝山钢铁股份有限公司
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Priority claimed from CN202110360571.6A external-priority patent/CN115181899B/zh
Priority claimed from CN202110360529.4A external-priority patent/CN115181896B/zh
Priority claimed from CN202110360520.3A external-priority patent/CN115181892B/zh
Priority claimed from CN202110360524.1A external-priority patent/CN115181893B/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
Publication of WO2022206912A1 publication Critical patent/WO2022206912A1/zh

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • 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
    • 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
    • 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

Definitions

  • the invention belongs to the technical field of rapid heat treatment of materials, and particularly relates to low-carbon low-alloy TRIP steel with tensile strength ⁇ 980 MPa or low-carbon low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980 MPa and a manufacturing method thereof.
  • TRIP transformation-induced plasticity
  • the mechanical properties of TRIP steel are determined by the volume fraction and strength of ferrite, bainite and austenite, as well as the microstructure and distribution of each phase, especially the stability of retained austenite against strain-induced martensitic transformation.
  • the heat treatment process of TRIP steel is mainly composed of two main stages: austenitizing annealing process and bainite isothermal treatment process.
  • the cold-rolled deformed matrix structure takes the lead in the recovery and recrystallization process.
  • the cementite in the matrix begins to dissolve in the ferrite.
  • the heating temperature exceeds A C1 , if the temperature is high and the time is sufficient, the cementite can be completely dissolved in the austenite to complete the austenite process.
  • critical annealing the carbon in the austenite phase is enriched to the A c3 line, after that, if the pro-eutectoid cementite is suppressed by alloying elements Si, Al, etc., the carbon concentration reaches T0 or T′ during austempering 0 .
  • bainite transformation begins to occur at the supercooled austenite grain boundary.
  • the carbon content is lower than the carbon content in austenite.
  • the austenite grain boundary forms bainite, the remaining carbon diffuses into the austenite that has not undergone the transformation reaction to form carbon-rich austenite.
  • the subsequent cooling process of the carbon-rich austenite stops the phase transformation, resulting in the formation of retained austenite.
  • the microstructure of TRIP steel is changed by adding alloying elements and adjusting the temperature and time during quenching and partitioning in the TRIP process. performance.
  • Chinese patent application CN102312157B discloses "a cold-rolled TRIP steel above 1000MPa grade and its preparation method", the chemical composition mass percentage of the invention steel is: C: 0.18-0.23%, Si: 1.3-1.64%, Mn: 2.1-2.3 %, Nb: 0.03 ⁇ 0.05%, V: 0.03 ⁇ 0.09%, P ⁇ 0.01%, S ⁇ 0.01%, Alt: 0.8 ⁇ 1.2%, N ⁇ 0.005%, the balance is Fe and other unavoidable impurity elements.
  • the main feature of the invention steel is that it is produced by using a traditional continuous production line, the soaking time is 3-8min; the heating rate is 1.5-15°C/s, the soaking temperature is 810-830°C, the soaking time is 85-120s, and the The holding time of intenite is 5-8min, and the annealing and bainite transformation treatment time are very long. At the same time, it adds higher micro-alloying elements such as C, Si, Mn, and Nb, V to replace Ni and Cr. In the actual production process, the alloy cost and the manufacturing cost in the production process are bound to be higher.
  • Chinese patent application CN109182923A discloses "a heat treatment method of low-carbon microalloyed high-strength plastic-deposited cold-rolled TRIP980 steel", the chemical composition mass percentage of the invention steel is: C: 0.18-0.23%, Si: 1.6-1.8%, Mn : 1.5 ⁇ 2.0%, Nb: 0.025 ⁇ 0.045%, Ti: 0.08 ⁇ 0.15%, P ⁇ 0.015%, S ⁇ 0.005%, the balance is Fe and other unavoidable impurity elements.
  • the main features of the invention steel are: on the one hand, a higher Si content is adopted, and higher microalloying elements Nb and Ti are added to obtain high elongation (A% ⁇ 24%) and high strength ( ⁇ 980MPa) .
  • the cold-rolled strip is subjected to two heat treatments.
  • the cold-rolled cold-rolled strip after pickling is first subjected to a complete austenitizing annealing and then quenching.
  • the whole martensitic structure is formed, and then the surface scale is removed and the decarburized layer is removed, and then heating and annealing is performed again to finally obtain the finished strip steel.
  • this method not only has the problems of high Si content and high addition of microalloying elements, but also the process is complicated, and two anneals are required, which leads to a significant increase in manufacturing cost and difficulty in manufacturing process.
  • Cide patent application 201711385126.5 discloses "a 780MPa low-carbon low-alloy TRIP steel", and its chemical composition mass percentage is: C: 0.16-0.22%, Si: 1.2-1.6%, Mn: 1.6-2.2%, and the balance is Fe and other unavoidable impurity elements, which are obtained by the following rapid heat treatment process: the strip steel is rapidly heated from room temperature to a two-phase region of austenite and ferrite at 790-830 °C, and the heating rate is 40-300 °C/s; The residence time in the target temperature range of heating in the two-phase zone is 60-100s; the strip steel is rapidly cooled from the temperature of the two-phase zone to 410-430°C, and the cooling rate is 40-100°C/s, and stays in this temperature range for 200-300s; The strip is rapidly cooled from 410 to 430°C to room temperature.
  • the metallographic structure of the TRIP steel is a three-phase structure of bainite, ferrite and austenite; the average grain size of the TRIP steel is obviously refined; the tensile strength is 950-1050MPa; 21 ⁇ 24%; the maximum strong-plastic product can reach 24GPa%.
  • the patent discloses a 780MPa grade low carbon and low alloy TRIP steel product and its process technology, but the tensile strength of the TRIP steel product is 950-1050MPa, which is too high as the tensile strength of a 780MPa grade product. It is impossible for the user to use the effect well, and the tensile strength of the 980MPa level is too low, which cannot well meet the user's strength requirements;
  • the patent adopts one-stage rapid heating, and the same rapid heating rate is used in the entire heating temperature range. It does not need to be treated differently according to the changes in the structure of the material in different temperature sections. Rapid heating will inevitably lead to an increase in the production cost of the rapid heating process;
  • the soaking time of the patent is set at 60-100s, which is similar to the soaking time of the traditional continuous retreat.
  • the increase of soaking time will inevitably partially weaken the grain refinement effect produced by rapid heating, which is very important for the improvement of material strength and toughness. unfavorable;
  • the patent must carry out the bainite isothermal treatment time of 200-300s, which is actually too long for the rapid heat treatment products, and the effect is limited and unnecessary. Moreover, the increase of soaking time and isothermal treatment time is not conducive to saving energy, reducing the investment in unit equipment and the floor space of the unit, and it is not conducive to the high-speed and stable operation of the strip in the furnace. Obviously, this is not a rapid heat treatment process in the strict sense. .
  • Chinese patent CN103237905B discloses "a multi-phase steel", a cold-rolled flat steel product made of this kind of multi-phase steel, and the manufacturing method.
  • the chemical composition mass percentage of the invention steel is: C: 0.14-0.25%, Mn: 1.7-2.5%, Si: 1.4-2.0%, Al ⁇ 0.1%, Cr ⁇ 0.1%, Mo ⁇ 0.05%, Nb: 0.02-0.06 %, S ⁇ 0.01%, P ⁇ 0.02%, N ⁇ 0.01%, and optionally at least one element from the "Ti, B, V" group: Ti ⁇ 0.1%, B ⁇ 0.002%, V ⁇ 0.15%,
  • the balance is Fe and other unavoidable impurity elements.
  • the hot rolling rolling temperature is 1100-1300 °C
  • the final rolling temperature is 820-950 °C
  • the coiling temperature is 400-750 °C; after 30-80% cold rolling, heat treatment is performed.
  • Over-aging treatment is carried out in the temperature range of 350-500°C.
  • High elongation (A%: 16.9-22%) and high strength (980-1086 MPa) were finally obtained.
  • the main feature of the invention steel is that it adopts high microalloying elements Nb, Mo, Cr, Ti or B, V (one must be included), and the content of its main alloy components C, Si and Mn is not low.
  • the traditional heat treatment process has the problems of low structure control ability, long annealing time and high cost.
  • Chinese patent application CN105543674B discloses "a manufacturing method of cold-rolled ultra-high-strength dual-phase steel with high local formability".
  • the chemical composition of the high-strength dual-phase steel of the invention is calculated as: C: 0.08-0.12%, Si: 0.1-0.5%, Mn: 1.5-2.5%, Al: 0.015-0.05%, and the rest are Fe and other inevitable impurities.
  • the chemical composition is matched with raw materials and smelted into a cast slab; the cast slab is heated at 1150-1250 °C for 1.5-2 hours and then hot rolled, the hot rolling rolling temperature is 1080-1150 °C, and the final rolling temperature is 880-930 °C; After rolling, it is cooled to 450-620°C at a cooling rate of 50-200°C/s for coiling to obtain a hot-rolled steel sheet with bainite as the main structure type; Heating to 740-820°C at a rate of /s for annealing, holding time for 30s-3min, cooling to 620-680°C at a cooling rate of 2-6°C/s, and then cooling to 250°C at a cooling rate of 30-100°C/s -350°C over-aging treatment for 3-5min to obtain ultra-high-strength dual-phase steel with ferrite + martensite dual-phase structure.
  • the ultra-high-strength dual-phase steel has a yield strength of 650-680 MPa, a tensile strength of 1023-1100 MPa, and an elongation of 12.3-13%. 180° bending along the rolling direction without cracking.
  • the most important feature of this patent is to combine the cooling condition control after hot rolling with the rapid heating in the continuous annealing process, that is, by controlling the cooling process after hot rolling, the strip structure is eliminated and the structure is homogenized; in the subsequent continuous annealing process Rapid heating is used to achieve tissue refinement on the basis of ensuring tissue uniformity. It can be seen that the patented technology adopts rapid heating and annealing. The premise is that hot-rolled raw materials with bainite as the main structure are obtained after hot-rolling.
  • the hot-rolled raw materials have high strength and high deformation resistance, which brings great difficulties to subsequent pickling and cold rolling production;
  • the soaking time is 30s-3min.
  • the increase of soaking time will inevitably partially weaken the grain refinement effect produced by rapid heating, which is not conducive to the improvement of material strength and toughness;
  • the patent must be over-aged for 3-5 minutes, which is actually too long for rapid heat treatment of DP steel and is not necessary. Moreover, the increase of soaking time and over-aging time is not conducive to saving energy, reducing the investment in unit equipment and the floor area of the unit, and it is not conducive to the high-speed and stable operation of the strip in the furnace. Obviously, this is not a rapid heat treatment process in the strict sense. .
  • Chinese patent application CN108774681A discloses "a rapid heat treatment method for high-strength steel".
  • the method uses a ceramic sheet electric heating device to obtain a heating rate with a maximum value of 400°C/s. After heating to 1000-1200°C, a fan is used to blow Cool down to room temperature with the fastest cooling rate of nearly 3000°C/s.
  • the processing speed of the heat treatment device using the electric heating of the ceramic sheet is 50 cm/min.
  • the steel of the invention is characterized in that its carbon content is as high as 0.16-0.55%, and simultaneously contains: Si, Mn, Cr, Mo and other alloy elements; the method is mainly suitable for steel wires, wire rods or steel strips below 5 mm.
  • the patent describes a rapid heat treatment method by electric heating of ceramic sheets; the main purpose of the invention is to solve the problems of low heat treatment efficiency, waste of energy and environmental pollution of products such as high-strength steel wires and wire rods; there is no mention of the effect of rapid heating on the material structure. Influence and function of performance; this invention does not combine the composition and structure characteristics of steel grades, and adopts fan blowing cooling method. At the same time, the use of excessive cooling rate in the high temperature section to produce wide and thin strips will lead to problems such as excessive internal stress and poor steel plate shape, which is not suitable for large-scale industrial continuous heat treatment production of wide and thin steel plates.
  • Chinese patent application CN107794357B and US patent application US2019/0153558A1 disclose "a method for producing ultra-high-strength martensitic cold-rolled steel sheet by ultra-rapid heating process", and the chemical composition of the high-strength steel is calculated as: C: 0.10 ⁇ 0.30%, Mn: 0.5-2.5%, Si: 0.05-0.3%, Mo: 0.05-0.3%, Ti: 0.01-0.04%, Cr: 0.10-0.3%, B: 0.001-0.004%, P ⁇ 0.02%, S ⁇ 0.02%, the rest is Fe and other inevitable impurities.
  • the invention provides an ultra-rapid heating production process for ultra-high strength martensitic cold-rolled steel sheets.
  • the heating rate of /s is heated to 850-950 °C in the single-phase austenite region; the steel plate is kept warm for no more than 5s and immediately water-cooled to room temperature to obtain an ultra-high-strength cold-rolled steel plate.
  • the annealing temperature of the steel of the invention has entered the ultra-high temperature range of the austenite single-phase region, and it also contains a lot of alloying elements, and the yield strength and tensile strength both exceed 1000MPa, so the heat treatment process, heat treatment The previous process and subsequent user use bring greater difficulties.
  • the ultra-rapid heating annealing method of the invention which adopts a holding time of no more than 5s, not only has poor controllability of the heating temperature, but also leads to uneven distribution of alloying elements in the final product, resulting in uneven product structure and performance. unstable;
  • the final quick cooling adopts water quenching to room temperature without necessary tempering treatment, so that the microstructure and properties of the final product and the distribution of alloying elements in the final microstructure cannot make the product obtain the best quality.
  • the strength and toughness of the final product is more than excess, but the plasticity and toughness are insufficient;
  • the method of the invention will cause problems such as poor shape and surface oxidation of the steel plate due to the high water quenching speed, so the patented technology has little or no practical application value.
  • the traditional continuous annealing production line also has a large number of furnace rolls in the high-temperature furnace section.
  • the traditional continuous annealing unit is based on the product outline and production capacity requirements Generally, the soaking time is required to be 1-3min.
  • the number of furnace rolls in the high-temperature furnace section varies from 20 to 40. This increases the difficulty of strip surface quality control.
  • Advanced high-strength steels represented by transformation-induced plasticity TRIP steel 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 and production of TRIP steel.
  • the purpose of the present invention is to provide a kind of low-carbon low-alloy TRIP steel with tensile strength ⁇ 980MPa and its rapid heat treatment manufacturing method, through rapid heat treatment to change the recovery, recrystallization and phase transformation process of the deformed structure, shorten the grain growth time, Refining the grains, the metallographic structure of the TRIP steel obtained is a multiphase structure of bainite, ferrite and austenite, with an average grain size of 1-3 ⁇ m; bainite is submicron granular; austenite The body is equiaxed grains with island-like distribution; bainite and austenite are evenly distributed on the ferrite matrix; its yield strength ⁇ 540Pa, tensile strength ⁇ 980MPa, elongation ⁇ 18%, strong-plastic product ⁇ 23GPa%; good plasticity and toughness are obtained while obviously improving the strength and performance of the material; the rapid heat treatment process improves the production efficiency, reduces the production cost and energy consumption, significantly reduces the
  • the technical scheme of the present invention is:
  • the content of C is selected from the range of 0.17-0.23%, 0.19-0.21%, 0.19-0.25% and 0.21-0.23%;
  • the content range is selected from: 1.1-1.7%, 1.3-1.5%, 1.3-2.0% and 1.5-1.8%; ⁇ 2.2%.
  • the content of Cr is ⁇ 0.3%
  • the content of Mo is ⁇ 0.3%
  • the content of Nb is ⁇ 0.05%
  • the content of Ti is ⁇ 0.04%
  • the content of V is ⁇ 0.04%. ⁇ 0.055%.
  • the low carbon and low alloy TRIP steel is obtained by the following process, including:
  • the final rolling temperature of hot rolling is ⁇ A r3 , and then it is cooled to 550 ⁇ 680°C for coiling;
  • the cold rolling reduction rate is 40-80%;
  • the cold-rolled steel sheet is rapidly heated to 770-860°C, and the rapid heating adopts a one-stage or two-stage type; when one-stage rapid heating is used, the heating rate is 50-500°C; when two-stage rapid heating is used, the first stage is used. Heating from room temperature to 550-625°C at a heating rate of 15-500°C/s, and heating the second stage from 550-625°C to 770-770°C at a heating rate of 30-500°C/s (eg 50-500°C/s). 860°C; then soaking, soaking temperature is 770-860°C, soaking time is 40-120s;
  • the isothermal treatment is carried out in the temperature range, and the isothermal treatment time is 150-250 s; after that, it is cooled to room temperature at a cooling rate of 30-150° C./s, preferably 30-100° C./s.
  • the entire process of the rapid heat treatment takes 281 to 350 s.
  • 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 550-625°C at a heating rate of 30-300°C/s; the second stage is heated at a heating rate of 50-300°C/s The heating rate is from 550 to 625 °C to 770 to 860 °C.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-625°C at a heating rate of 50-300°C/s; the second stage is heated at a heating rate of 80-300°C/s The heating rate is from 550 to 625 °C to 770 to 860 °C.
  • the chemical composition mass percentage of the low-carbon low-alloy TRIP steel with a tensile strength ⁇ 980 MPa of the present invention is: C: 0.17-0.23%, Si: 1.1-1.7%, Mn: 1.6-2.2%, P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, can also contain one or two of Cr, Mo, Ti, Nb, V, Cr+Mo+Ti+Nb+V ⁇ 0.5%, The balance is Fe and other inevitable impurities.
  • the C content is 0.19-0.21%.
  • the Si content is 1.3-1.5%.
  • the Mn content is 1.8-2.0%.
  • the metallographic structure of the low-carbon low-alloy TRIP steel is a multiphase structure of 35-45% bainite, 10-60% ferrite, and 5-15% austenite.
  • the low-carbon low-alloy TRIP steel has a yield strength ⁇ 540 MPa, a tensile strength ⁇ 980 MPa, an elongation ⁇ 21%, and a strong-plastic product ⁇ 23 GPa%; more preferably, the low-carbon low-alloy TRIP steel has a yield strength of ⁇ 980 MPa It is 549-716MPa, the tensile strength is 1030-1120MPa, the elongation is 21.3-24.5%, and the strong-plastic product is 23-25.9GPa%.
  • the chemical composition mass percentage of the low-carbon low-alloy TRIP steel with tensile strength ⁇ 980 MPa according to the present invention is: C: 0.19-0.25%, Si: 1.3-2.0%, Mn: 1.8-2.4%, P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, can also contain one or two of Cr, Mo, Ti, Nb, V, Cr+Mo+Ti+Nb+V ⁇ 0.5%, the rest The amount is Fe and other inevitable impurities.
  • the tensile strength of the low-carbon low-alloy TRIP steel is ⁇ 1180 MPa.
  • the C content is 0.21-0.23%.
  • the Si content is 1.5-1.8%.
  • the Mn content is 2.0-2.2%.
  • the metallographic structure of the low-carbon low-alloy TRIP steel is a three-phase structure of 40-80% bainite, 10-50% ferrite, and 7-18% austenite.
  • the yield strength of the low-carbon low-alloy TRIP steel is greater than or equal to 770MPa, the tensile strength is greater than or equal to 1180MPa, the elongation is greater than or equal to 17%, and the strength-plastic product is greater than or equal to 23GPa%; 861MPa, the tensile strength is 1180-1297MPa, the elongation is 17-21%, and the strong-plastic product is 23-25.9GPa%.
  • Another aspect of the present invention provides a low-carbon and low-alloy hot-dip galvanized TRIP steel with a tensile strength of ⁇ 980 MPa and a method for manufacturing hot-dip galvanizing by rapid heat treatment. Grain growth time, grain refinement, the metallographic structure of TRIP steel is obtained.
  • Bainite is submicron granular; austenite is equiaxed grains with island-like distribution; bainite and austenite are uniformly distributed
  • the yield strength is ⁇ 540MPa
  • the tensile strength is ⁇ 980MPa
  • the elongation is ⁇ 18%
  • the strong-plastic product is ⁇ 20.5GPa%
  • the rapid heat treatment process improves production efficiency, reduces production costs and energy consumption, significantly reduces the number of furnace rolls, and improves the surface quality of the steel strip.
  • the chemical composition mass percentage of the low-carbon low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980MPa is: C: 0.17-0.25%, Si: 1.1-2.0%, Mn: 1.6-2.4%, P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, can also contain one or two of Cr, Mo, Ti, Nb, V, Cr+Mo+Ti+Nb+V ⁇ 0.5%, the balance is Fe and other inevitable impurities.
  • the metallographic structure of the low-carbon low-alloy hot-dip galvanized TRIP steel is a three-phase structure of bainite, ferrite and austenite, and the average grain size is 1-3 ⁇ m; preferably, the bainite is Submicron granular, austenite is equiaxed grains with island-like distribution, and bainite and austenite are uniformly distributed on the ferrite matrix; preferably, in the multiphase structure, the volume of bainite is The volume ratio of ferrite is 35-80%, such as 35-75% or 40-80%, the volume ratio of ferrite is 10-60%, such as 10-50%, and the volume ratio of austenite is 5-18%, such as 7- 18% or 5-15%; preferably, the low-carbon low-alloy TRIP steel has a yield strength of ⁇ 540 MPa, a tensile strength of ⁇ 980 MPa, an elongation of ⁇ 18%, and a strong-plastic product of ⁇ 23
  • the content of C is selected from the range of 0.17-0.23%, 0.19-0.21%, 0.19-0.25% and 0.21-0.23%; preferably , the content range of Si is selected from: 1.1-1.7%, 1.3-1.5%, 1.3-2.0%, 1.5-1.8% and 1.5-1.9%; preferably, the content of Mn is selected from the range of 1.6-2.2%, 1.8-2.0% %, 1.8 to 2.4% and 2.0 to 2.2%.
  • the content of Cr is ⁇ 0.3%
  • the content of Mo is ⁇ 0.3%
  • the content of Nb is ⁇ 0.05%
  • the content of Ti is ⁇ 0.04%
  • the content of V is ⁇ 0.055%.
  • the low-carbon and low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980 MPa is obtained by the following process:
  • the final rolling temperature of hot rolling is ⁇ A r3 , and then it is cooled to 550 ⁇ 680°C for coiling;
  • the cold rolling reduction rate is 40-80%;
  • the cold-rolled steel plate is rapidly heated to 770-860°C, and the rapid heating adopts a one-stage or two-stage type;
  • the heating rate is 50-500°C/s;
  • the first stage is heated from room temperature to 550-625°C at a heating rate of 15-500°C/s, and the second stage is heated at a heating rate of 50-500°C/s (such as 50-500°C/s). Heating rate from 550 ⁇ 625°C to 770 ⁇ 860°C; then soaking, soaking temperature: 770 ⁇ 860°C, soaking time: 30 ⁇ 120s;
  • the isothermal treatment time is 60-150s; after the isothermal treatment is completed, the heating is heated to 460-470°C at a heating rate of 10-30°C/s, and then immersed in a zinc pot for hot-dip galvanizing;
  • hot-dip galvanizing After hot-dip galvanizing, rapidly cool to room temperature at a cooling rate of 30-150°C/s to obtain hot-dip pure zinc GI products; ) at a heating rate of 480-550°C for alloying treatment, and the alloying treatment time is 5-20s; after alloying treatment, it is rapidly cooled to room temperature at a cooling rate of 30-250°C/s (eg 30-100°C/s). , to obtain alloyed hot-dip galvanized GA products.
  • the whole process of rapid heat treatment and hot-dip galvanizing takes 118-328s, such as 118-238s.
  • 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 550-625°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 to 625 °C to 770 to 860 °C.
  • the rapid heating adopts two-stage heating: the first stage is heated from room temperature to 550-625°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 625 °C to 770 to 860 °C.
  • the final temperature of the rapid heating is 790-860°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 860°C, and the temperature after cooling is not lower than 770°C.
  • step 4 after the strip or steel plate is hot-dip galvanized, it is heated to 480-550°C at a heating rate of 30-200°C/s for alloying treatment, and the alloying treatment time is 5-20s; After treatment, it is rapidly cooled to room temperature at a cooling rate of 30-200° C./s to obtain alloyed hot-dip galvanized GA products.
  • the chemical composition mass percentage of the low-carbon low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980 MPa is: C: 0.17-0.23%, Si: 1.1-1.7%, Mn: 1.6-2.2% , P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, 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 is 0.19-0.21%.
  • the Si content is 1.3-1.5%.
  • the Mn content is 1.8-2.0%.
  • the whole process of the rapid heat treatment and hot-dip galvanizing takes 118-328s.
  • the metallographic structure of the low-carbon low-alloy hot-dip galvanized TRIP steel is a three-phase structure of 35-75% bainite, 10-60% ferrite, and 5-15% austenite.
  • the yield strength of the low-carbon low-alloy hot-dip galvanized TRIP steel is 549-620 MPa
  • the tensile strength is increased to 1030-1164 MPa
  • the elongation is 20.1-24.4%
  • the strong-plastic product is 20.7-25.8 GPa%.
  • the chemical composition mass percentage of the low-carbon low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980 MPa is: C: 0.19-0.25%, Si: 1.3-2.0%, Mn: 1.8-2.4% , P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, 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 is 0.21-0.23%.
  • the Si content is 1.5-1.9%.
  • the Mn content is 2.0-2.2%.
  • the whole process of the rapid heat treatment and hot-dip galvanizing takes 118-238s.
  • the metallographic structure of the low-carbon low-alloy hot-dip galvanized TRIP steel is a three-phase structure of 40-80% bainite, 10-50% ferrite, and 7-18% austenite.
  • the yield strength of the low-carbon low-alloy hot-dip galvanized TRIP steel is 771-821 MPa
  • the tensile strength is increased to 1182-1284 MPa
  • the elongation is 18-22.2%
  • the strong-plastic product is 22.6-26.4 GPa%.
  • 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 generally exists in steel in several ways: solid solution carbon and cementite in ferrite and austenite. Carbon content has a great influence on the mechanical properties of steel. With the increase of carbon content, the number of strengthening phases such as bainite, pearlite and martensite will increase, which will greatly improve the strength and hardness of steel, but its plasticity and toughness will be significantly reduced. If the carbon content is too 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. The effect will also be significantly weakened, so that the process performance of the steel will be deteriorated, so the carbon content should be reduced as much as possible under the premise of ensuring the strength.
  • the solid solution of carbon in the austenite can expand the austenite phase region, increase the amount of retained austenite, improve its stability, and make the transformation of pearlite and bainite easier.
  • the C curve shifts to the right, delaying the transformation of ferrite and bainite and reducing the Ms point temperature.
  • the content of carbon in the austenite determines the amount and degree of stability of the retained austenite. The higher the carbon content of the retained austenite, the better the stability of the retained austenite. As the carbon content increases, the content of retained austenite also increases.
  • the present invention limits the carbon content within the range of 0.17-0.25%. In some embodiments, the C content is 0.17-0.23%. In other embodiments, the C content is 0.19-0.25%.
  • 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 during the cooling process after hot rolling. The content will also increase with the increase of Mn content.
  • Manganese is also a carbide forming element. 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.
  • Si and Mn the presence of Si element will aggravate the segregation degree of Mn element, strengthen the dragging effect of Mn on C atoms, and delay the formation of bainite.
  • the manganese content is high, the structure will be banded and the residual austenite will be too stable, which is not conducive to the occurrence of phase transformation; it will also cause the grains in the steel to coarsen and increase the overheating sensitivity of the steel. Improper cooling after rolling can easily cause white spots in carbon steel.
  • the present invention controls the manganese content within the range of 1.6-2.4%. In some embodiments, the Mn content is 1.6-2.2%. In some embodiments, the Mn content is 1.8-2.4%.
  • Si Silicon forms a solid solution in ferrite or austenite, thereby enhancing the yield strength and tensile strength of steel. Silicon can increase the cold working deformation hardening rate of steel and is a beneficial element in alloy steel. In addition, 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, and it will not significantly reduce the plasticity of steel within a certain range. Silicon has a strong deoxidizing 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. However, if the content of silicon in steel is too high, its plasticity and toughness will decrease significantly.
  • Si element is a ferrite-forming element, which can improve the stability of retained austenite, and also play a role in solid solution strengthening to improve the strength of the steel. Silicon also has the effect of reducing the austenite phase region and increasing the activity of C element in ferrite. Higher silicon content is beneficial to obtain more retained austenite, but too high silicon content will cause steel to produce such as hard oxide layer, poor surface properties, reduced wettability of hot-rolled steel sheets, and deteriorated surface quality, etc. question. If the content of silicon is too low, a stable and satisfactory TRIP effect will not be brought, so the content of silicon must be controlled within a certain range.
  • Silicon has no significant effect on the growth rate of austenite, but has a significant effect on the morphology and distribution of the formed austenite.
  • the increase of the silicon content will increase the difficulty of manufacturing in the process before the heat treatment. Based on the above factors, the present invention controls the silicon content within the range of 1.1-2.0%. In some embodiments, the Si content is 1.1-1.7%. In some embodiments, the Si content is 1.3-2.0%.
  • Chromium and iron form a continuous solid solution, which narrows the austenite phase region. Chromium and carbon form various carbides, and their 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 is added with other alloying elements in improving the strength and hardness of steel, and 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 content of Cr is ⁇ 0.3%.
  • 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 strong carbide forming elements V, Nb and Ti are added to the steel, the solid solution strengthening effect of Mo is more significant.
  • 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 content of Mo is ⁇ 0.3%.
  • Micro-alloying elements Ti, Nb, V adding a small amount of micro-alloying elements Nb, V, Ti to the steel can ensure that the steel can pass through its carbon and nitride material points (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 strong carbide and nitride forming elements, and these elements can meet this requirement at relatively low concentrations. At room temperature, most of them exist in the form of carbides, nitrides, and carbonitrides in steel, and a small part is solid-dissolved in ferrite.
  • the addition of microalloying elements can strengthen the ferrite matrix through grain refinement and precipitation, and can also delay the formation of bainite.
  • the reason for the delayed bainite formation is the intensified ferrite formation upon cooling, which is the result of microstructural grain refinement.
  • the formation of ferrite leads to carbon enrichment of retained austenite, delaying the transformation of austenite to bainite, while finely dispersed carbonitrides inhibit bainite nucleation, which also delays bainite formation.
  • the addition of Nb, V and Ti can prevent the growth of austenite grains and increase the roughening temperature of the steel. This is because their small particles dispersed in carbon and nitride can fix the austenite grain boundaries and hinder the austenite grain boundaries.
  • the migration of body grain boundaries increases the recrystallization temperature of austenite, which can expand the unrecrystallized area, that is, prevent the growth of austenite grains.
  • 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 content of Nb is ⁇ 0.05%. In some embodiments, the content of Ti is ⁇ 0.04%. In some embodiments, the content of V is ⁇ 0.055%.
  • the invention finely controls the recovery, recrystallization and phase transformation process of the deformed structure of the hard-rolled strip steel during the heat treatment process through the rapid heat treatment method (including rapid heating, short-term heat preservation and rapid cooling process), and finally obtains fine, uniform and dispersed distribution.
  • the rapid heat treatment method including rapid heating, short-term heat preservation and rapid cooling process
  • 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 or grow significantly.
  • the heat treatment process disclosed in Chinese patent application CN106811698B does not differentiate the entire heating process, and the heating rate used in the heating process is 20-60°C/s, which is a medium heating rate, and is based on the heating of the existing traditional continuous annealing unit. It is not possible to carry out large-scale regulation according to the needs of material organization transformation.
  • the multiphase structure of ferrite, austenite and bainite obtained by the rapid heat treatment method of the present invention has an average grain size of 1-3 ⁇ m, which is smaller than the grain size of products produced by the existing traditional technology (usually 50 to 80% at 5 to 15 ⁇ m), the strength of the material can be improved through grain refinement, and good plasticity and toughness can be obtained at the same time, and the performance of the material can be improved; and the ferrite, bainite and The retained austenite structure is mainly in various forms such as massive and granular, and the distribution is more uniform, so that better strong plasticity can be obtained in the deformation stage.
  • the rapid heat treatment manufacturing method of low-carbon and low-alloy TRIP steel with tensile strength ⁇ 980 MPa comprises the following steps:
  • the final rolling temperature of hot rolling is ⁇ A r3 , and then it is cooled to 550 ⁇ 680°C for coiling;
  • the cold rolling reduction rate is 40% to 80%, and the rolled hard strip or steel plate is obtained;
  • the strip or steel plate after cold rolling is rapidly heated to 770 ⁇ 860°C, and the rapid heating adopts one-stage or two-stage type; when one-stage rapid heating is adopted, the heating rate is 50 ⁇ 500°C/s; two-stage rapid heating is adopted When the first stage is heated from room temperature to 550-625°C at a heating rate of 15-500°C/s, the second stage is heated from 550-625°C at a heating rate of 30-500°C/s (eg 50-500°C/s). °C heated to 770 ⁇ 860 °C;
  • Soaking is carried out at the target temperature of austenite and ferrite two-phase region at 770-860 °C, and the soaking time is 40-120s;
  • Bainite isothermal treatment is carried out in the temperature range of 410 ⁇ 430 °C, and the isothermal treatment time is 150 ⁇ 250s;
  • the strip or steel plate is cooled to room temperature at a cooling rate of 30-150°C/s.
  • the entire process of the rapid heat treatment takes 281 to 350 s.
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the final temperature of the rapid heating is 790-830°C.
  • the heating rate is 50-300°C/s.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-625°C at a heating rate of 30-300°C/s, and the second stage is heated at a heating rate of 50-300°C/s
  • the heating rate is from 550 to 625 °C to 770 to 860 °C.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-625°C at a heating rate of 50-300°C/s, and the second stage is heated at a heating rate of 80-300°C/s
  • the heating rate is from 550 to 625 °C to 770 to 860 °C.
  • the cooling rate of the rapid cooling process is 50-100°C/s.
  • 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.
  • step 4) strip steel or steel plate is carried out small-amplitude heating or small-amplitude cooling in soaking time section, and the temperature after heating is not more than 860 °C, and the temperature after cooling is not less than 770 °C.
  • the rapid heat treatment manufacturing method of low-carbon low-alloy hot-dip galvanized TRIP steel with tensile strength ⁇ 980 MPa comprises the following steps:
  • the final rolling temperature of hot rolling is ⁇ A r3 , and then it is cooled to 550 ⁇ 680°C for coiling;
  • the cold rolling reduction rate is 40 to 80%, and the rolled hard strip or steel plate is obtained after cold rolling;
  • the heating rate is 50-500°C/s;
  • the first stage is heated from room temperature to 550-625°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 ⁇ 625°C to 770 ⁇ 860°C;
  • Soaking is carried out at a target temperature of 770-860°C in the two-phase region of austenite and ferrite, and the soaking time is 30-120s;
  • 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 410-430°C at a cooling rate of 40-100°C/s;
  • the strip or steel plate is subjected to bainite isothermal treatment at 410 ⁇ 430°C, and the isothermal treatment time is 60 ⁇ 150s;
  • the strip or steel plate is then dipped into a zinc pot for hot-dip galvanizing;
  • the entire process of the rapid heat treatment in step 4) takes 118-328 s, such as 118-238 s.
  • 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 550-625°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 to 625 °C to 770 to 860 °C.
  • the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-625°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 625 °C to 770 to 860 °C.
  • the final temperature of the rapid heating is 790-860°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 860°C, and the temperature after cooling is not lower than 770°C.
  • step 4 after the strip or steel plate is hot-dip galvanized, it is heated to 480-550°C at a heating rate of 30-200°C/s for alloying treatment, and the alloying treatment time is 5-20s; After treatment, it is rapidly cooled to room temperature at a cooling rate of 30-200° C./s to obtain alloyed hot-dip galvanized GA products.
  • the recrystallization kinetics of the continuous heating process is 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 is limited by the influence of heating technology, which is 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, and the phase deformation occurs.
  • the nucleation is mainly at the ferrite grain boundaries that have grown, and the nucleation rate is low. Therefore, the resulting tissue is relatively coarse.
  • the deformed matrix has just completed recrystallization or has not completed recrystallization (even not fully recovered), and the phase transformation of ferrite to austenite begins to occur, because recrystallization has just completed or has not been completed.
  • 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.
  • the ferrite recrystallization and the austenite transformation process overlap. Since a large number of crystal defects such as dislocations remain in the ferrite crystal, a large number of nucleation points are provided for the formation of austenite, which makes austenite austenite.
  • the nucleation shows explosive nucleation, so the austenite grains are further refined.
  • 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 quickly grow and grow, so the austenite grains are fine and the volume fraction increases.
  • the present invention sets the heating rate as 50-500°C/s during one-stage rapid heating, and sets the heating rate as 15 ⁇ 500°C/s.
  • the optimal heating rate in different heating temperature ranges are also different: from 20°C to 500-625°C, the heating rate has the greatest influence on the recovery process, and the heating rate is controlled to be 15-300°C/s, more preferably 30-300°C/s; the heating temperature is from 500-625°C When the austenitizing temperature is 770-860°C, the heating rate has the greatest influence on the grain growth process, and the heating rate is controlled to be 50-300°C/s; more preferably 80-300°C/s.
  • the choice of soaking temperature needs to be combined with the control of the evolution of the material at each temperature stage of the heating process, and at the same time, the evolution and control of the structure in the subsequent rapid cooling process must be considered, so that the optimal structure and distribution can be finally obtained.
  • the soaking temperature usually depends on the C content in the steel. In the traditional heat treatment process, the soaking temperature is generally set between A C1 and A C3 , or 30 to 50°C above A c3 .
  • the invention uses the rapid heating technology to retain a large number of crystal defects such as dislocations in the matrix that is not fully recovered and recrystallized, and provides sufficient nucleation work for austenite transformation, so the temperature only needs to be heated to A C1 to A C3 In between, more austenite can be obtained.
  • the C content of the TRIP steel is: 0.17-0.25%, and A C1 and A C3 are respectively about 730°C and about 870°C.
  • a small amount of fine granular undissolved carbides should be evenly distributed in the alloy 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.
  • the selection of soaking temperature should also aim to obtain fine and uniform austenite grains, so as to obtain fine and uniform ferrite, bainite and retained austenite after cooling. Excessive soaking temperature will make the austenite grains coarse, the workpiece is easy to crack during the quenching process, and the microstructure obtained after quenching will also be coarser, resulting in poor mechanical properties of the steel. If the soaking temperature is too low, the content of carbon and alloying elements dissolved in the austenite will be insufficient, resulting in uneven distribution of the austenite carbon concentration, which will greatly reduce the hardenability of the steel, and will be detrimental to the mechanical properties of the alloy steel. influences.
  • the soaking temperature of hypoeutectoid steel should be Ac3+30 ⁇ 50°C.
  • the present invention selects 770-860° 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. When the content of carbon and alloying elements in the steel increases, it will not only reduce the thermal conductivity of the steel, but also because the diffusion rate of alloying elements is faster than that of carbon elements. Slowly, the alloying elements will obviously delay the structural transformation of the steel, and the holding time should be appropriately extended at this time.
  • the material in the two-phase region contains a large number of dislocations, provides a large number of nucleation points for the formation of austenite, and provides a rapid diffusion channel for carbon atoms, so the austenite can be formed extremely quickly, and
  • the shorter the soaking and holding time the shorter the diffusion distance of carbon atoms, the greater the carbon concentration gradient in the austenite, and the more retained austenite carbon content at the end; however, if the holding time is too short, the distribution of alloying elements in the steel will be increased. Uneven, resulting in insufficient austenitization; too long holding time will easily lead to coarse austenite grains.
  • the influencing factors of soaking and holding time also depend on the content of carbon and other alloying elements in the steel.
  • the control of soaking time needs to be formulated in strict combination with soaking temperature, rapid cooling and rapid heating process, in order to finally obtain the ideal structure and element distribution.
  • the soaking time is set as 40-120s.
  • the control of the rapid cooling process needs to combine the comprehensive factors such as the results of the evolution of the structures and the results of the alloy diffusion distribution in the pre-heating and soaking processes to ensure that the ideal phase structure and the material structure with reasonable distribution of elements are finally obtained.
  • the cooling rate of the material during rapid cooling must be greater than the critical cooling rate to obtain bainite.
  • the critical cooling rate mainly depends on the material composition.
  • the Si content in the TRIP steel of the present invention is 1.1% to 2.0%, and the Mn content It is 1.6-2.4%, and the content is relatively high, so Si and Mn greatly enhance the hardenability of austenite in TRIP steel and reduce the critical cooling rate.
  • the cooling rate also needs to comprehensively consider the microstructure evolution of the heating process and the soaking process and the alloy diffusion distribution results, so as to finally obtain a reasonable microstructure distribution of each phase and alloy element distribution. If the cooling rate is too low, the bainite structure cannot be obtained, and the mechanical properties cannot meet the requirements. However, if the cooling rate is too large, a large quenching stress (ie, organizational stress and thermal stress) will be generated, which will cause poor shape and even lead to testing. Severely deformed and cracked. Therefore, in the present invention, the rapid cooling rate is set to 40-100°C/s.
  • the bainite isothermal treatment temperature of TRIP steel is generally selected below the temperature (T 0 ) equal to the free energy of bainite, ferrite and austenite. At this time, the free energy of bainite is less than the free energy of austenite, and the The reduction provides the phase transformation driving force for the bainite transformation. Due to the different chemical components of steel materials, the bainite isothermal treatment temperature is also different, and the bainite isothermal temperature is generally selected between 350 and 550 °C. The isothermal treatment temperature is high, the atomic diffusion ability is strong, the austenite is partially transformed into granular bainite, and carbides are precipitated, which reduces the stability of the supercooled austenite, and the residual austenite volume fraction is low.
  • bainite transformation that requires atomic diffusion is difficult to carry out, which may lead to martensite transformation without atomic diffusion.
  • Martensite is a supersaturated structure of C.
  • C diffusion is too slow. , it is difficult to enrich in supercooled austenite, and it will also lead to the reduction of the volume fraction of retained austenite, so the isothermal temperature of bainite in the present invention is selected between 410 and 430 °C.
  • the bainite isothermal treatment time is short, the bainite transformation is not fully carried out, the enrichment degree of C element to austenite is low, and the austenite C content is low, so its stability is poor, and in the subsequent cooling process In the process, the supercooled austenite will be transformed into martensite, and the martensite structure has the characteristics of high strength and low elongation, so it is not good for improving the strong plasticity.
  • the bainite transformation is sufficient, and the bainite volume fraction in the TRIP steel of the present invention increases.
  • the isothermal time prolongs, the SEM microstructure does not change significantly, and the bainite volume fraction and morphology change little.
  • the present invention sets the bainite isothermal time to be longer, at 60-250s, such as 60-150s or 150-250s.
  • the rapid heat treatment process reduces the residence time of the strip in the high-temperature furnace, so the enrichment of alloying elements on the surface of the high-strength strip is significantly reduced during the heat treatment process, which is conducive to improving the High-strength hot-dip galvanized products can be plated, reduce surface leakage plating defects, and improve corrosion resistance, thereby increasing the yield.
  • the present invention can greatly shorten the length of the heating and soaking section of the annealing furnace by transforming the traditional continuous annealing unit to the rapid heating and rapid cooling process to realize the rapid heat treatment process (at least one-third shorter than the traditional continuous annealing furnace). 1), improve the production efficiency of the traditional continuous annealing unit, reduce the production cost and energy consumption, and significantly reduce the number of continuous annealing furnace rolls, especially the number of high-temperature furnace rolls, which can improve the strip surface quality control ability and obtain high Surface quality strip products.
  • 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 by rapid heat treatment, so that the recrystallization process overlaps with the austenite transformation process, and the recrystallization grains and austenite grains are increased. nucleation point, shorten the grain growth time, and refine the grains.
  • the microstructure of the obtained TRIP steel is that the volume of bainite structure accounts for 35-80%, the volume of ferrite structure accounts for 15-60%, The multiphase structure of the austenite structure accounts for 5-18% of the volume, and the average grain size is refined to 1-3 ⁇ m, which is 50-50% smaller than the grain size (usually 5-15 ⁇ m) of the products produced by the existing traditional technology.
  • the microstructure of the obtained hot-dip galvanized TRIP steel is a three-phase structure of bainite (35-80%), ferrite (10-60%) and austenite (5-18%).
  • the average grain size is 1 to 3 ⁇ m, and the average grain size is reduced by 30 to 50%.
  • bainite is submicron granular; austenite is island-shaped grains; bainite and austenite are evenly distributed on the ferrite matrix; austenite also has good thermal properties. Stability, -50°C austenite transformation rate is lower than 8%; -190°C austenite transformation rate is lower than 30%, and austenite of different sizes, shapes and orientations can continue to TRIP under different strain conditions effect and significantly improve the material properties.
  • the strength of the material can be improved, while good plasticity and toughness can be obtained, and the performance of the material can be improved.
  • the TRIP steel obtained by the invention has a multi-phase structure, and the grain size is reduced by 50-80%, the strength and toughness of the material are obviously improved, and its tensile strength can be controlled at a relatively low level.
  • the small range such as 1030-1120MPa, or 1190-1300MPa
  • the stability of the mechanical properties of the product is significantly improved, and the elongation remains at a high level (such as 21.3-24.5%).
  • the TRIP steel obtained by the present invention has a multi-phase structure with fine grains, the average grain size is 1-3 ⁇ m, and the average grain size is reduced by 30-50%; It can obviously improve the strength and toughness of the material, the yield strength is 549-821MPa, the tensile strength is increased to 1030-1284MPa; the elongation is 18-24.4%; the strong-plastic product is 20.7-26.4GPa%.
  • the time for the whole heat treatment process can be shortened to 281-350s;
  • the time of the whole heat treatment process is greatly reduced (the traditional continuous annealing process time of TRIP steel is usually 9-11min); especially the residence time at high temperature above 600°C is shortened, thereby improving production efficiency, reducing energy consumption and reducing production cost.
  • the rapid heat treatment method of the present invention shortens the heating section and soaking section time by 60-80%, and shortens the treatment time of the strip at high temperature, and the entire heat treatment process.
  • the time can be shortened to 281-350s, which can save energy, reduce emissions, reduce consumption, significantly reduce the one-time investment in furnace equipment, significantly reduce production and operation costs and equipment maintenance costs; in addition, the production of products of the same strength grade through rapid heat treatment can reduce the alloy content and reduce heat treatment. And the production cost of the previous process, reducing the manufacturing difficulty of each process before the heat treatment.
  • the low-carbon and low-alloy TRIP steel with a tensile strength of ⁇ 980 MPa or the low-carbon and low-alloy hot-dip galvanized TRIP steel with a tensile strength of ⁇ 980 MPa can be used for the development of a new generation of lightweight vehicles, trains, ships, airplanes and other transportation vehicles. It is of great value to the healthy development of the corresponding industries and advanced manufacturing industries.
  • Fig. 1 is the microstructure picture of TRIP steel produced according to Example 1 of Test Steel A in Example 1 of the present invention.
  • FIG. 2 is a picture of the microstructure of the TRIP steel produced by the traditional process 1 of the test steel A in Example 1 of the present invention.
  • Example 3 is a picture of the microstructure of the TRIP steel produced according to Example 6 of the test steel P in Example 1 of the present invention.
  • Example 4 is a picture of the microstructure of the TRIP steel produced according to Example 12 of the test steel M in Example 1 of the present invention.
  • FIG. 5 is a picture of the microstructure of the TRIP steel produced according to Example 21 of Test Steel G in Example 1 of the present invention.
  • FIG. 6 is a picture of the microstructure of the TRIP steel produced according to Example 23 of the test steel S in Example 1 of the present invention.
  • Example 7 is a picture of the microstructure of the TRIP steel produced in Example 1 of the second test steel A of the present invention.
  • Fig. 8 is the microstructure picture of the TRIP steel produced by the traditional process 1 of the test steel A in the second embodiment of the present invention.
  • Example 9 is a picture of the microstructure of the TRIP steel produced in Example 6 of the second test steel P of the present invention.
  • Example 10 is a picture of the microstructure of the TRIP steel produced in Example 12 of the second test steel M of the present invention.
  • Example 11 is a picture of the microstructure of the TRIP steel produced in Example 21 of Test Steel G in Example 2 of the present invention.
  • FIG. 12 is a picture of the microstructure of the TRIP steel produced in Example 23 of the test steel S in Example 2 of the present invention.
  • Example 13 is a microstructure picture of the hot-dip pure zinc TRIP steel (GI) produced according to Example 1 of Test Steel A in Example 3 of the present invention.
  • Example 15 is a microstructure picture of the alloyed hot-dip galvanized TRIP steel (GA) produced in Example 17 of Test Steel I in Example 3 of the present invention.
  • Example 16 is a microstructure picture of the hot-dip pure zinc TRIP steel (GI) produced according to Example 22 of Test Steel D in Example 3 of the present invention.
  • Example 17 is a microstructure picture of the alloyed hot-dip galvanized TRIP steel (GA) produced according to Example 34 of Test Steel I in Example 3 of the present invention.
  • Example 18 is a microstructure picture of the hot-dip pure zinc TRIP steel (GI) produced in Example 1 of the test steel A in Example 4 of the present invention.
  • FIG. 20 is a microstructure picture of the alloyed hot-dip galvanized TRIP steel (GA) produced according to Example 17 of Test Steel I in Example 4 of the present invention.
  • FIG. 21 is a microstructure picture of the hot-dip pure zinc TRIP steel (GI) produced according to Example 22 of Test Steel D in Example 4 of the present invention.
  • Example 22 is a microstructure picture of the alloyed hot-dip galvanized TRIP steel (GA) produced by Example 4 of the present invention, Test Steel I according to Example 34.
  • the yield strength, tensile strength and elongation were carried out according to "GB/T228.1-2010 Metal Materials Tensile Test Part 1: Test Method at Room Temperature", and the P7 sample was used to test in the transverse direction.
  • the alloy content in the same grade of steel can be reduced, the grains can be refined, and a good match of the material structure and strength and toughness can be obtained.
  • the yield strength of the TRIP steel obtained by the method of the present invention is 549-716 MPa, the tensile strength is 1030-1120 MPa, the elongation is 21.3-24.5%, and the strong-plastic product is 23-25.9 GPa%.
  • FIG. 1 is a microstructure diagram of a typical composition A steel obtained by Example 1
  • FIG. 2 is a microstructure diagram of a typical composition A steel obtained by a traditional process Example 1. From the figure, there are very big differences in the microstructure after different heat treatment methods.
  • the microstructure (Fig. 1) of steel A is mainly composed of uniform bainite and retained austenite dispersed and distributed on the ferrite matrix.
  • the structure treated by the process of the invention: the bainite and the retained austenite structure are very uniformly distributed in the ferrite matrix, which is very beneficial for improving the strength and plasticity of the material.
  • the structure of A steel treated by the traditional process (Fig. 2) is a typical TRIP steel structure.
  • Bainite and a small amount of retained austenite are distributed on the bulk ferrite grain boundaries, and the average grain size is about 5-15 ⁇ m.
  • the characteristics of the structure treated by the traditional process are: the grain structure is relatively coarse, the content of bainite and retained austenite structure is relatively small, and the distribution is relatively uneven.
  • FIG. 3 is a microstructure diagram of typical composition P steel obtained by Example 6, and FIG. 4 is a microstructure diagram of typical composition M steel obtained by Example 12.
  • FIG. 5 is a microstructure diagram of typical composition G steel obtained by Example 21, and
  • FIG. 6 is a microstructure diagram of typical composition S steel obtained by Example 23.
  • Examples 6, 12, 21, 23 are typical rapid heat treatment processes. It can be seen from the figure that, by using the method of the present invention, very uniform, fine, and dispersed ferrite, bainite and retained austenite phase structures can be obtained.
  • the preparation method of the TRIP 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 method of the present invention can reduce the alloy content in the steel of the same grade, refine the grains, and obtain a good match of the material structure and strength and toughness.
  • the yield strength of the TRIP steel obtained by the method of the present invention is 751-947 MPa
  • the tensile strength is 1193-1297 MPa
  • the elongation is 18.5-21%
  • the strong-plastic product is 23-25.9 GPa%.
  • FIG. 7 is a microstructure diagram of a typical composition A steel obtained by Example 1
  • FIG. 8 is a microstructure diagram of a typical composition A steel obtained by a traditional process Example 1. From the figure, there are very big differences in the microstructure after different heat treatment methods.
  • the structure of the TRIP steel obtained after the treatment in the embodiment of the present invention is mainly composed of ferrite, bainite, retained austenite and a small amount of carbides, and the bainite and retained austenite are very uniformly distributed in the iron In the element body matrix, this is very beneficial to improve the strength and plasticity of the material.
  • the structure of TRIP steel obtained by traditional processing is a typical TRIP steel structure.
  • Bainite and a small amount of retained austenite are distributed on the grain boundary of bulk ferrite, and the average grain size is about 5-15 ⁇ m. .
  • the characteristics of the structure treated by the traditional process are: the grain structure is relatively coarse, the content of bainite and retained austenite structure is relatively small, and the distribution is relatively uneven.
  • FIG. 9 is a microstructure diagram of typical composition P steel obtained by Example 6, and FIG. 10 is a microstructure diagram obtained by typical composition M steel of Example 12.
  • FIG. 11 is a microstructure diagram of typical composition G steel obtained by Example 21, and
  • FIG. 12 is a microstructure diagram of typical composition S steel obtained by Example 23.
  • Examples 6, 12, 21, 23 are typical rapid heat treatment processes. It can be seen from the figure that, by using the method of the present invention, very uniform, fine, and dispersed ferrite, bainite and retained austenite phase structures can be obtained.
  • the preparation method of the TRIP 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, so as to realize the rapid heat treatment process, greatly shorten the length of the heating section and soaking section of the traditional continuous annealing furnace, and improve the
  • the production efficiency of the traditional continuous annealing unit can reduce the production cost and energy consumption, and reduce the number of furnace rolls in the continuous annealing furnace, which can improve the control ability of the strip surface quality and obtain strip products with high surface quality; at the same time, by establishing a rapid heat treatment process
  • the new continuous annealing unit of technology makes the continuous heat treatment unit short and compact, flexible in material transition, and strong in control ability; for the material, it can refine the grain of the strip steel, further improve the material strength, reduce the cost of alloy and the manufacturing process before heat treatment Difficulty, improve the welding performance of materials 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 tens of seconds, tens of 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. For the unit speed of 180 meters The number of rollers in the high temperature furnace section of the rapid heat treatment production line of about /min can be no 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 period of time will also provide a more flexible and flexible high-strength steel structure design method, without changing the alloy composition and rolling process.
  • the material structure can be improved and the material properties can be improved.
  • Advanced high-strength steels represented by transformation-induced plasticity TRIP steel have broad application prospects, and rapid heat treatment technology has great development and application value. The combination of the two will surely provide more space for the development and production of TRIP steel. .
  • Table 11 The composition of the test steel of the present invention is shown in Table 11, and the specific parameters of the present embodiment and the traditional process are shown in Table 12 (one-stage heating) and Table 13 (two-stage heating); Table 14 and Table 15 are the composition of the test steel of the present invention according to the table Main properties of GI and GA hot-dip galvanized TRIP steels prepared by the examples and conventional processes in Table 12 and Table 13.
  • 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 TRIP steel obtained by the method of the invention can reach 549-620MPa, the tensile strength is increased to 1030-1164MPa, the elongation is 20.1-24.4%, and the strong-plastic product is 20.7-25.8GPa%.
  • Fig. 13 and Fig. 14 are the structure diagrams of typical composition A steel through Example 1 and Comparative Traditional Process Example 1. From the two figures, there is a very big difference in the structure of the two processes after hot-dip galvanizing.
  • the microstructure (Fig. 13) of the steel A after the rapid heat treatment of the present invention is mainly composed of fine and uniform bainite structure and carbides dispersed on the fine ferrite matrix, and the bainite grain structure and carbide composition All are very fine and uniformly dispersed, which is very beneficial to improve the strength and plasticity of the material.
  • the structure of A steel treated by the traditional process (Fig. 14) is a typical TRIP steel structure.
  • the structure of the material treated by the traditional process shows a certain directionality, and the structure is elongated along the rolling direction.
  • the characteristics of the structure treated by the traditional process are: the grain size is large, and there is a certain band structure, the bainite and retained austenite are distributed in a network along the ferrite grain boundary, the ferrite grain is relatively coarse, the ferrite The two-phase microstructure distribution of body and bainite is uneven.
  • Fig. 15 is a microstructure diagram of typical composition I steel obtained by Example 17 (GA)
  • Fig. 16 is a microstructure diagram of a typical composition D steel obtained by Example 22 (GI).
  • Figure 17 is a microstructure diagram of a typical composition I steel obtained through Example 34 (GA).
  • Embodiments 17, 22, and 34 are all processes with a shorter entire heat treatment cycle.
  • the rapid heat treatment hot-dip galvanizing method of the present invention a very uniform, fine and dispersed phase structure is obtained after alloying treatment, while in the metallographic structure of the steel strip prepared by the traditional process, ferrite is The structure is coarse, and the bainite and retained austenite structure is distributed on the ferrite grain boundary, which is a typical hot-dip galvanized TRIP steel structure. Therefore, the preparation method of the hot-dip galvanized TRIP 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.
  • Table 16 for the composition of the test steel in Example 4, and Table 17 (one-stage heating) and Table 18 (two-stage heating) for the specific parameters of the embodiment of the present invention and the traditional process;
  • Table 19 and Table 20 are the test steels of the present invention
  • Table 19 is the test steel composition of the present invention according to the embodiment and the traditional process in Table 17 and Table 18.
  • the obtained GI and GA hot-dip galvanized TRIP steel are prepared. main performance.
  • 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 TRIP steel obtained by the method of the invention can reach 771-821MPa, the tensile strength is increased to 1182-1284MPa, the elongation is 18-22.2%, and the strong-plastic product is 22.6-26.4GPa%.
  • Fig. 18 and Fig. 19 are the structure diagrams of typical composition A steel through Example 1 and Comparative Traditional Process Example 1. From the two figures, there is a very big difference in the structure of the two processes after hot-dip galvanizing.
  • the microstructure of the steel A after the rapid heat treatment of the present invention (Fig. 18) is mainly composed of microstructures such as ferrite, bainite and retained austenite which are dispersed and distributed uniformly. Various structures are very fine and uniformly dispersed, which is very beneficial to improve the strength and plasticity of the material.
  • the structure of A steel treated by the traditional process is a typical TRIP steel structure.
  • the characteristics of the structure treated by the traditional process are: the structure is distributed in a certain direction along the rolling direction.
  • the proportion of ferrite structure is slightly larger, bainite and retained austenite structure are distributed along the rolling direction, and the structure has certain inhomogeneity.
  • Fig. 20 is a microstructure diagram of a typical composition I steel obtained by Example 17 (GA)
  • Fig. 21 is a microstructure diagram of a typical composition D steel obtained by Example 22 (GI).
  • Figure 22 is a microstructure diagram of a typical composition I steel obtained through Example 34 (GA).
  • Embodiments 17, 22, and 34 are all processes with a shorter entire heat treatment cycle. As can be seen from the figure, using the rapid heat treatment hot-dip galvanizing method of the present invention, very uniform, fine, and dispersed phase structures are obtained after alloying treatment (Fig. 20), while the steel strip metallographic structure prepared by the traditional process 9 is obtained. Among them, the microstructures such as ferrite and bainite are relatively coarse and have certain directionality.
  • the preparation method of the hot-dip galvanized TRIP 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.
  • Examples 3 and 4 show that by adopting the rapid heating and rapid cooling process to transform the traditional continuous annealing hot-dip galvanizing unit to realize the rapid heat-treatment hot-dip galvanizing process, the traditional continuous annealing hot-dip galvanizing process can be greatly shortened.
  • the length of the heating section and soaking section of the furnace can improve the production efficiency of the traditional continuous annealing hot-dip galvanizing unit, reduce the production cost and energy consumption, reduce the number of furnace rolls in the continuous annealing hot-dip galvanizing furnace, and significantly reduce the roll marks, pitting, and rubbing.
  • the hot-dip galvanizing unit can be realized. It has the advantages of being short and compact, flexible in material transition, and strong in control ability; for hot-dip galvanized substrate materials, it can refine grains, further improve material strength, reduce alloy costs and the difficulty of manufacturing before heat treatment, and improve material forming, welding, etc. User performance.
  • the present invention greatly promotes the technological progress of the continuous annealing hot-dip galvanizing process of the cold-rolled strip steel.
  • the integration process can be expected to be completed within ten seconds or even a few seconds, which greatly shortens the length of the heating section of the continuous annealing hot-dip galvanizing furnace, facilitates the improvement of the speed and production efficiency of the continuous annealing hot-dip galvanizing unit, and significantly reduces the continuous annealing hot-dip galvanizing unit.
  • the rapid heat treatment hot-dip galvanizing process of recrystallization and austenitization completed in a very short time will also provide a more flexible and flexible high-strength steel microstructure design method, and then do not need to change the alloy composition and rolling. Under the premise of pre-process conditions such as technology, the structure of the material can be improved and the performance of the material can be improved.
  • the hot-dip galvanized advanced high-strength steel represented by the transformation-induced plasticity TRIP steel has broad application prospects, and the rapid heat treatment technology has great development and application value. Production provides more space.

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Abstract

抗拉强度≥980MPa的低碳低合金TRIP钢或低碳低合金热镀锌TRIP钢，其化学成分质量百分比为：C：0.17～0.25％，Si：1.1～2.0％，Mn：1.6～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。其制造方法包括：冶炼、铸造、热轧、冷轧和快速热处理或快速热处理热镀；快速热处理全过程用时281～350s，快速热处理热镀全过程用时118～328s。本发明通过控制快速热处理过程中快速加热、短时保温和快速冷却过程，改变变形组织的回复、再结晶及奥氏体相变过程，增加形核率，缩短晶粒长大时间，细化晶粒，最终获得TRIP钢显微组织为贝氏体、铁素体、奥氏体多相组织，且平均晶粒尺寸在1～3μm；在提高材料强度的同时获得良好的塑性和韧性。

Description

抗拉强度≥980MPa的低碳低合金TRIP钢或热镀锌TRIP钢及其制造方法 技术领域
本发明属于材料快速热处理技术领域，特别涉及抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢及其制造方法。
背景技术
随着人们对能源节约以及材料安全服役意识的逐步提高，高强钢，尤其是先进高强钢的使用日益增多，使得钢铁企业及科研院所对先进高强钢的开发日益重视。
为了进一步提高钢材产品的强塑积，特别是其延伸率，以TRIP(相变诱导塑性)钢为代表的先进高强钢的开发日益得到重视。冷轧相变诱导塑性TRIP钢组织是由铁素体基体上包含贝氏体和残余奥氏体构成，在塑性变形时，亚稳态残余奥氏体转变成为马氏体，因而在加工高强度时具有优良的伸长率和可成形性。通过亚稳态的残余奥氏体在很大应变范围内产生的形变诱导相变作用改善钢的成型性能，其总伸长率可达到30-40％，由于存在较软的多边形铁素体组织的原因，导致其抗拉强度一般不超过980MPa。TRIP钢的力学性能由铁素体、贝氏体和奥氏体的体积分数、强度及各相组织形态、分布决定，特别是残余奥氏体抑制应变诱导马氏体转变的稳定性。TRIP钢的热处理工艺过程，主要由奥氏体化退火过程和贝氏体等温处理过程两个主要阶段组成。
1.加热及奥氏体化过程
在连续加热退火过程中，冷轧变形基体组织率先发生回复与再结晶过程，在该温度区间，基体中的渗碳体开始溶解于铁素体中。当加热温度超过A C1之后，如果温度高且时间充裕，渗碳体可全部溶解于奥氏体中，完成奥氏体化过程。通过临界退火，奥氏体相中的碳富集到A c3线，此后，如果先共析渗碳体被合金元素Si、Al等抑制，则在等温淬火过程中碳浓度达到T 0或T′ 0。
2.快速冷却及贝氏体等温过程
材料完成奥氏体化之后进行快速冷却，当冷却到贝氏体相变温度时进行保温(等温)处理，首先在过冷奥氏体晶界处开始发生贝氏体转变，由于贝氏体中的碳含量低于奥氏体中的碳含量，奥氏体晶界形成贝氏体后，剩余碳向尚未发生相变反应的奥氏体中扩散形成富碳奥氏体，当富碳奥氏体中的碳含量达到某一临界值时，富碳奥氏体随后冷却过程就会 停止相转变，从而形成残余奥氏体。
目前，针对相变诱导塑性TRIP钢成分配方及生产工艺的研究开发、优化完善，主要是通过添加合金元素、调整TRIP工艺中淬火及配分过程中温度及时间来改变相变诱导塑性TRIP钢的组织性能。
中国专利申请CN102312157B公开了“一种1000MPa级以上冷轧TRIP钢及其制备方法”，该发明钢化学成分质量百分比为：C：0.18～0.23％，Si：1.3～1.64％，Mn：2.1～2.3％，Nb：0.03～0.05％，V：0.03～0.09％，P≤0.01％，S≤0.01％，Alt：0.8～1.2％，N≤0.005％，余量为Fe和其它不可避免的杂质元素。该发明钢的主要特征在于采用传统连退产线进行生产，均热时间3-8min；其加热速率为1.5-15℃/s，均热温度810-830℃，均热时间85-120s，贝氏体保温时间5-8min，其退火及贝氏体转变处理时间都很长。同时其添加较高C、Si、Mn及Nb、V等微合金元素，替代Ni、Cr，实际生产过程其合金成本和生产工艺过程中的制造成本必然也较高。
中国专利申请CN109182923A公开了“一种低碳微合金化高强塑积冷轧TRIP980钢的热处理方法”，该发明钢化学成分质量百分比为：C：0.18～0.23％，Si：1.6～1.8％，Mn：1.5～2.0％，Nb：0.025～0.045％，Ti：0.08～0.15％，P≤0.015％，S≤0.005％，余量为Fe和其它不可避免的杂质元素。该发明钢的主要特征在于：一方面，采用了较高的Si含量，同时添加较高的微合金元素Nb、Ti，来获得高延伸率(A％≥24％)和高强度(≥980MPa)。另一方面，与传统TRIP钢生产工艺相比，其采用的是对冷轧带钢进行两次热处理，酸洗后冷轧处理的冷轧带钢首先进行一次完全奥氏体化退火，然后淬火成全马氏体组织，随后进行表面除氧化铁皮和去除脱碳层，再重新进行一次加热退火，最终得到成品带钢。显然该方法不仅存在Si含量偏高和微合金元素添加量偏高问题，而且工艺复杂，要采用两次退火，导致制造成本显著增加、制造工序难度增加。
中国专利申请201711385126.5公开了“一种780MPa级别低碳低合金TRIP钢”，其化学成分质量百分比为：C：0.16-0.22％，Si：1.2-1.6％，Mn：1.6-2.2％，余量为Fe和其它不可避免的杂质元素，其通过下述快速热处理工艺获得：带钢由室温快速加热至790～830℃奥氏体和铁素体两相区，加热速率为40～300℃/s；在两相区加热目标温度区间停留时间为60～100s；带钢从两相区温度快速冷却至410～430℃，冷却速度为40～100℃/s，并在此温度区间停留200-300s；带钢从410～430℃快速冷却至室温。其特征在于：所述的TRIP钢金相组织为贝氏体、铁素体、奥氏体三相组织；所述的TRIP钢平均晶粒尺寸明显细化；抗拉强度950～1050MPa；延伸率21～24％；强塑积最大可达到24GPa％。
该专利的不足主要有以下几个方面：
第一，该专利公开的是一种780MPa级别低碳低合金TRIP钢产品及其工艺技术，但该TRIP钢产品的抗拉强度为950～1050MPa该强度作为780MPa级的产品抗拉强度显得太高了，用户使用效果不可能好，而作为980MPa级别抗拉强度又偏低了，不能很好地满足用户的强度要求；
第二，该专利采用一段式快速加热，在整个加热温度区间均采用了同一个快速加热速率，未根据不同温度段的材料组织结构变化需要进行区别处理，而全部以40～300℃/s的速度快速加热，必然导致快速加热过程生产成本的提高；
第三，该专利均热时间定为60～100s，这和传统连退的均热时间差不多，均热时间的增加必然部分减弱快速加热产生的细化晶粒效果，对材料强度和韧性提高非常不利；
第四，该专利必须进行200～300s的贝氏体等温处理时间，这实际上对快速热处理产品而言等温处理时间过长了，起到的作用有限，没有必要。而且均热时间和等温处理时间的增加都不利于节约能源、降低机组设备投资和机组占地面积，更不利于带钢在炉内的高速稳定运行，显然这也不是严格意义上的快速热处理过程。
中国专利CN103237905B公开了“一种多相钢”，由该类多相钢制成的、冷轧的扁钢制品以及该制造方法。该发明钢化学成分质量百分比为：C：0.14～0.25％，Mn：1.7～2.5％，Si：1.4～2.0％，Al≤0.1％，Cr≤0.1％，Mo≤0.05％，Nb：0.02～0.06％，S≤0.01％，P≤0.02％，N≤0.01％，以及可选的“Ti、B、V”族中的至少一个元素：Ti≤0.1％，B≤0.002％，V≤0.15％，余量为Fe和其它不可避免的杂质元素。其热轧开轧温度1100-1300℃，终轧温度820-950℃，卷取温度400-750℃；经过30-80％冷轧压下后进行热处理，热处理采用传统退火工艺方式进行，退火温度A C1+20℃～A C3范围内。在350-500℃温度范围内进行过时效处理。最终获得高延伸率(A％：16.9-22％)和高强度(980-1086MPa)。该发明钢的主要特征在于采用添加较高微合金元素Nb、Mo、Cr、Ti或B、V(必须含有之一)，其主合金成分C、Si、Mn的含量也不低。同时采用传统热处理工艺，存在组织控制能力低、退火时间长，成本高等问题。
中国专利申请CN105543674B公开了“一种高局部成型性能冷轧超高强双相钢的制造方法”，该发明的高强度双相钢化学成分按重量百分数计为：C：0.08～0.12％、Si：0.1～0.5％、Mn：1.5～2.5％、Al：0.015～0.05％，其余为Fe和其它不可避免杂质。将该化学成分选配原料，熔炼成铸坯；将铸坯在1150-1250℃加热1.5-2小时后进行热轧，热轧开轧温度1080-1150℃，终轧温度为880-930℃；轧后以50-200℃/s的冷却速度冷却至450-620℃进行卷取，得到以贝氏体为主要组织类型的热轧钢板；将热轧钢板进行冷轧，随后以50-300℃/s的速度加热至740-820℃进行退火，保温时间30s-3min，以2-6℃/s的冷速冷至 620-680℃，之后以30-100℃/s的冷速冷至250-350℃过时效处理3-5min，得到铁素体+马氏体双相组织的超高强双相钢。该超高强双相钢的屈服强度为650-680MPa，抗拉强度为1023-1100MPa，延伸率为12.3-13％。沿轧制方向180°弯曲不开裂。
该专利的最主要特征为将热轧后冷却条件控制与连续退火过程中的快速加热相结合，即通过控制热轧后冷却工艺，消除带状组织，实现组织均匀化；在后续连续退火过程中采用快速加热，在保证组织均匀性的基础上实现组织细化。可见该专利技术采用快速加热退火，其前提是热轧后获得以贝氏体为主要组织的热轧原料，其目的主要在于保证组织均匀性，避免出现带状组织而导致局部变形不足。
该专利的不足之处在于：
第一，要获得具有贝氏体组织的热轧原料，该热轧原料强度高、变形抗力大，为后续酸洗和冷轧生产都带来了很大的困难；
第二，其对快速加热的理解仅限于缩短加热时间，细化晶粒的层面，其加热速率未根据不同温度段的材料组织结构变化需要进行划分，而全部以50-300℃/s的速度加热，导致快速加热生产成本的提高；
第三，均热时间30s-3min，均热时间的增加必然部分减弱快速加热产生的细化晶粒效果，对材料强度和韧性提高不利；
第四，该专利必须进行3-5分钟的过时效处理，这实际上对快速热处理DP钢而言时效时间过长了，没有必要。而且均热时间和过时效时间的增加都不利于节约能源、降低机组设备投资和机组占地面积，更不利于带钢在炉内的高速稳定运行，显然这也不是严格意义上的快速热处理过程。
中国专利申请CN108774681A公开了“一种高强钢的快速热处理方法”，该方法采用陶瓷片电加热装置，可获得最大值达到400℃/s的加热速率，加热到1000～1200℃后，采用风机吹风冷却，最快冷速近3000℃/s的冷速冷至室温。该发明方法中采用陶瓷片电加热的热处理装置处理速度为50cm/min。该发明钢的特征在于其含碳量高达0.16～0.55％，且同时含有：Si、Mn、Cr、Mo等合金元素；该方法主要适合于钢丝、盘条或5mm以下的钢带。该专利阐述了一种通过陶瓷片电加热的快速热处理方法；该发明的主要目的在于解决高强钢丝和盘条等产品热处理效率低，浪费能源及环境污染的问题；未提及快速加热对材料组织性能的影响及作用；该发明未结合钢种牌号成分及组织特点，采用风机吹风冷却的方式，最快冷速接近3000℃/s指的是高温段的瞬时冷速，平均冷速是达不到3000℃/s的；同时高温段采用过高的冷速生产宽薄带钢会导致内应力过大、钢板板型不良等问题，不适用于宽薄钢板的大规模工业化连续热处理生产。
中国专利申请CN107794357B和美国专利申请US2019/0153558A1公开了“一种超快速加热工艺生产超高强度马氏体冷轧钢板的方法”，该高强度钢化学成分按重量百分数计为：C：0.10～0.30％、Mn：0.5～2.5％、Si：0.05～0.3％、Mo：0.05～0.3％、Ti：0.01～0.04％、Cr：0.10～0.3％、B：0.001～0.004％、P≤0.02％、S≤0.02％，其余为Fe和其它不可避免杂质。该钢的力学性能：屈服强度R p0.2大于1100MPa，抗拉强度R m＝1800-2300MPa，延伸率最大12.3％，均匀延伸率5.5-6％。该发明提供了一种超高强度马氏体冷轧钢板的超快速加热生产工艺，其工艺特征首先将冷轧钢板以1～10℃/s加热到300～500℃，然后以100～500℃/s的加热速率加热至单相奥氏体区850～950℃；钢板保温不超过5s立即水冷到室温，得到超高强度冷轧钢板。
该专利所述工艺的不足之处包括：
第一，该发明钢退火温度已经进入到奥氏体单相区的超高温温度范围，而且还含有较多的合金元素，屈服强度和抗拉强度均超过了1000MPa，所以给本热处理工艺、热处理前工序及后续用户使用带来较大的困难。
第二，该发明的超快速加热退火方法，其采用不超过5s的保温时间，不仅加热温度的可控性差，而且还会导致最终产品中合金元素分布不均匀，导致产品组织性能的不均匀和不稳定；
第三，最后的快冷其采用的是水淬冷却到室温，未进行必要的回火处理，这样其所得到的最终产品组织性能及最终组织结构中的合金元素分布情况不能使产品获得最佳的强韧性，导致最终产品强度过剩有余，而塑性和韧性不足；
第四，该发明的方法由于水淬冷速过高会导致钢板板型不良和表面氧化等问题，因此该专利技术没有太大实际应用价值或实际应用价值不大。
以往受企业生产设备所限，绝大部分的相关研究都是基于在现有传统加热装备的加热速率(5～20℃/s)对带钢进行慢速加热，使其依次完成再结晶和奥氏体化相变，因此加热和均热时间都比较长、能耗高，同时传统连续退火生产线还存在带钢在高温炉段的炉辊数目较多，传统连续退火机组根据产品大纲和产能要求，一般均热时间要求在1-3min，对于机组速度在180米/分左右的传统产线其高温炉段内的炉辊数目在20-40根不等。这对带钢表面质量控制难度增大。
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在几十秒、十几秒、甚至几秒内完成奥氏体化过程，大大缩短了加热段长度，便于提高机组速度和生产效率。同时，极短时间内所完成的奥氏体化过程也将提供更加灵活及柔性化的组织设计和产线设计方法，进而在无需改变合金成 分以及轧制工艺的前提下改善TRIP钢材料性能。
以相变诱导塑性TRIP钢为代表的先进高强钢有着广阔的应用前景，而快速热处理技术有着巨大的开发价值，两者的结合必将会为TRIP钢的开发和生产提供更大的空间。
发明内容
本发明的目的在于提供一种抗拉强度≥980MPa的低碳低合金TRIP钢及其快速热处理制造方法，通过快速热处理改变变形组织的回复、再结晶及相变过程，缩短晶粒长大时间，细化晶粒，获得TRIP钢的金相组织为贝氏体、铁素体和奥氏体的多相组织，平均晶粒尺寸为1～3μm；贝氏体为亚微米级颗粒状；奥氏体为孤岛状分布的等轴晶粒；贝氏体和奥氏体均匀分布在铁素体基体上；其屈服强度≥540Pa，抗拉强度≥980MPa，延伸率为≥18％，强塑积≥23GPa％；在明显提高材料强度性能的同时获得良好的塑性和韧性；采用快速热处理工艺提高了生产效率，降低生产成本及能耗，显著减少炉辊数量，提高钢带表面质量。
为达到上述目的，本发明的技术方案是：
抗拉强度≥980MPa的低碳低合金TRIP钢，其化学成分质量百分比为：C：0.17～0.25％，Si：1.1～2.0％，Mn：1.6～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；优选地，该低碳低合金TRIP钢的金相组织为贝氏体、铁素体和奥氏体多相组织，平均晶粒尺寸为1～3μm；优选地，贝氏体为亚微米级颗粒状，奥氏体为孤岛状分布的等轴晶粒，贝氏体和奥氏体均匀分布在铁素体基体上；优选地，所述多相组织中，贝氏体的体积比为35～80％、如35～45％或40～80％，铁素体的体积比10～60％、如10～50％，奥氏体的体积比为5～18％、如7～18％或5～15％；优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率为≥18％，强塑积≥23GPa％。优选地，该低碳低合金TRIP钢的奥氏体具有良好的热稳定性，-50℃奥氏体转化变率低于8％，-190℃奥氏体转变率低于30％。
优选地，所述抗拉强度≥980MPa的低碳低合金TRIP钢中，C的含量范围选自0.17～0.23％、0.19～0.21％、0.19～0.25％和0.21～0.23％；优选地，Si的含量范围选自：1.1～1.7％、1.3～1.5％、1.3～2.0％和1.5～1.8％；优选地，Mn的含量范围选自1.6～2.2％、1.8～2.0％、1.8～2.4％和2.0～2.2％。
优选地，所述抗拉强度≥980MPa的低碳低合金TRIP钢中，Cr的含量≤0.3％，Mo的含量≤0.3％，Nb的含量≤0.05％，Ti的含量≤0.04％，V的含量≤0.055％。
优选地，该低碳低合金TRIP钢通过下述工艺获得，包括：
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
热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
3)冷轧
冷轧压下率为40～80％；
4)快速热处理
冷轧后的钢板快速加热至770～860℃，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℃；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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；之后进行均热，均热温度为770～860℃，均热时间为40～120s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以40～100℃/s的冷却速率快速冷却至贝氏体等温处理温度410～430℃，并在此温度区间进行等温处理，等温处理时间150～250s；之后以30～150℃/s、优选30～100℃/s的冷却速率冷却至室温。
优选的，所述快速热处理全过程用时为281～350s。
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30～300℃/s的加热速率从室温加热至550～625℃；第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃。
优选的，步骤4)中，所述快速加热采用两段式加热，第一段以50～300℃/s的加热速率从室温加热至550～625℃；第二段以80～300℃/s的加热速率从550～625℃加热至770～860℃。
在一些实施方案中，本发明所述抗拉强度≥980MPa的低碳低合金TRIP钢的化学成分质量百分比为：C：0.17～0.23％，Si：1.1～1.7％，Mn：1.6～2.2％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，该低碳低合金TRIP钢中，C含量为0.19～0.21％。优选地，该低碳低合金TRIP钢中，Si含量为1.3～1.5％。优选地，该低碳低合金TRIP钢中，Mn含量为1.8～2.0％。优选地，所述低碳低合金TRIP钢的金相组织为贝氏体35～45％、 铁素体10～60％、奥氏体5～15％多相组织。优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率≥21％，强塑积≥23GPa％；更优选地，所述低碳低合金TRIP钢的屈服强度为549～716MPa，抗拉强度为1030～1120MPa，延伸率为21.3～24.5％，强塑积23～25.9GPa％。
在一些实施方案中，本发明所述抗拉强度≥980MPa的低碳低合金TRIP钢化学成分质量百分比为：C：0.19～0.25％，Si：1.3～2.0％，Mn：1.8～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，该低碳低合金TRIP钢的抗拉强度≥1180MPa。优选地，该低碳低合金TRIP钢中，C含量为0.21～0.23％。优选地，该低碳低合金TRIP钢中，Si含量为1.5～1.8％。优选地，该低碳低合金TRIP钢中，Mn含量为2.0～2.2％。优选地，该低碳低合金TRIP钢的金相组织为贝氏体40～80％、铁素体10～50％、奥氏体7～18％的三相组织。优选地，该低碳低合金TRIP钢的屈服强度≥770MPa，抗拉强度≥1180MPa，延伸率≥17％，强塑积≥23GPa％；更优选地，该低碳低合金TRIP钢的为776～861MPa，抗拉强度为1180～1297MPa，延伸率为17～21％，强塑积为23～25.9GPa％。
本发明另一方面提供一种抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢及快速热处理热镀锌制造方法，通过快速热处理改变变形组织的回复、再结晶及相变过程，缩短晶粒长大时间，细化晶粒，获得TRIP钢的金相组织中贝氏体为亚微米级颗粒状；奥氏体为孤岛状分布的等轴晶粒；贝氏体和奥氏体均匀分布在铁素体基体上，其屈服强度≥540MPa，抗拉强度≥980MPa，延伸率≥18％，强塑积≥20.5GPa％；在明显提高材料的强度的同时获得良好的塑性和韧性；同时，采用快速热处理工艺提高了生产效率，降低生产成本及能耗，显著减少炉辊数量，提高钢带表面质量。
该抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的化学成分质量百分比为：C：0.17～0.25％，Si：1.1～2.0％，Mn：1.6～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，该低碳低合金热镀锌TRIP钢的金相组织为贝氏体、铁素体和奥氏体的三相组织，平均晶粒尺寸为1～3μm；优选地，贝氏体为亚微米级颗粒状，奥氏体为孤岛状分布的等轴晶粒，贝氏体和奥氏体均匀分布在铁素体基体上；优选地，所述多相组织中，贝氏体的体积比为35～80％、如35～75％或40～80％，铁素体的体积比10～60％、如10～50％，奥氏体的体积比为5～18％、如7～18％或5～15％；优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率为≥18％，强塑积≥23GPa％。优选地，该低碳低合金TRIP钢的奥氏体具有良好的热稳定性，-50℃奥氏体转化变率低于8％， -190℃奥氏体转变率低于30％。
优选地，所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢中，C的含量范围选自0.17～0.23％、0.19～0.21％、0.19～0.25％和0.21～0.23％；优选地，Si的含量范围选自：1.1～1.7％、1.3～1.5％、1.3～2.0％、1.5～1.8％和1.5～1.9％；优选地，Mn的含量范围选自1.6～2.2％、1.8～2.0％、1.8～2.4％和2.0～2.2％。
优选地，所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢中，Cr的含量≤0.3％，Mo的含量≤0.3％，Nb的含量≤0.05％，Ti的含量≤0.04％，V的含量≤0.055％。
优选地，该抗拉强度≥980MPa的低碳低合金热镀锌TRIP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
热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
3)冷轧
冷轧压下率为40～80％；
4)快速热处理、热镀锌
冷轧后的钢板快速加热至770～860℃，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550～625℃，第二段以50～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；之后进行均热，均热温度：770～860℃，均热时间：30～120s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以40～100℃/s的冷却速率快速冷却至410～430℃；并在此温度区间进行等温处理，等温处理时间60～150s；等温处理结束后以10～30℃/s的加热速率加热至460～470℃，随后浸入锌锅进行热镀锌；
热镀锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀纯锌GI产品；或者，热镀锌之后，以30～300℃/s(如30～100℃/s)的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～250℃/s(如30～100℃/s)的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
优选地，所述快速热处理、热镀锌全过程用时为118～328s，如118～238s。
优选地，步骤2)中，所述卷取温度为580～650℃。
优选地，步骤3)中，所述冷轧压下率为60～80％。
优选地，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。
优选地，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550～625℃，第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃。
优选地，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550～625℃，第二段以80～300℃/s的加热速率从550～625℃加热至770～860℃。
优选地，步骤4)中，所述快速加热最终温度为790～860℃。
优选地，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。
优选地，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860℃，降温后温度不低于770℃。
优选地，步骤4)中，所述带钢或钢板热镀锌之后，以30～200℃/s的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～200℃/s的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
在一些实施方案中，所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的化学成分质量百分比为：C：0.17～0.23％，Si：1.1～1.7％，Mn：1.6～2.2％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，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，所述低碳低合金热镀锌TRIP钢中，C含量为0.19～0.21％。优选地，所述低碳低合金热镀锌TRIP钢中，Si含量为1.3～1.5％。优选地，所述低碳低合金热镀锌TRIP钢中，Mn含量为1.8～2.0％。优选地，所述快速热处理、热镀锌全过程用时为118～328s。优选地，所述低碳低合金热镀锌TRIP钢的金相组织为贝氏体35～75％、铁素体10～60％、奥氏体5～15％的三相组织。优选地，所述低碳低合金热镀锌TRIP钢的屈服强度549～620MPa，抗拉强度提高至1030～1164MPa；延伸率20.1～24.4％；强塑积20.7～25.8GPa％。
在一些实施方案中，所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的化学成分质量百分比为：C：0.19～0.25％，Si：1.3～2.0％，Mn：1.8～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，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，所述低碳低合金热镀锌TRIP钢中，C含量为0.21～0.23％。优选地，所述低碳低合金热镀锌TRIP钢中，Si含量为1.5～1.9％。优选地，所述低碳低合金热镀锌TRIP钢中，Mn含量为2.0～2.2％。优选地，所述快速热处理、热镀锌全过程用时为118～238s。优选地，所述低碳低合金热镀锌TRIP钢的金相组 织为贝氏体40～80％、铁素体10～50％、奥氏体7～18％的三相组织。优选地，所述低碳低合金热镀锌TRIP钢的屈服强度为771～821MPa，抗拉强度提高为1182～1284MPa；延伸率为18～22.2％；强塑积为22.6～26.4GPa％。
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。碳在钢中一般有几种存在方式：铁素体和奥氏体中的固溶碳和渗碳体。碳含量对钢的力学性能影响十分大，随着含碳量的升高，贝氏体、珠光体及马氏体等强化相的数量会增加，使钢的强度与硬度大幅提高，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，从而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。
对于相变诱导塑性TRIP钢而言，碳元素固溶于奥氏体中，可以扩大奥氏体相区，增加残余奥氏体的数量，提高其稳定性，使珠光体和贝氏体转变的C曲线右移，推迟铁素体和贝氏体的转变，降低Ms点温度。奥氏体中碳的含量决定了残余奥氏体的量和稳定程度，残余奥氏体含碳量越高残余奥氏体的稳定性越好。随碳含量增加，残余奥氏体的含量也增加。但含碳量过高则会降低钢的焊接性能；含碳量太低则使残余奥氏体的稳定性大大降低，甚至没有TRIP效应出现。因此，本发明将含碳量限定在0.17～0.25％范围之内。在一些实施方案中，C含量为0.17～0.23％。在另外一些实施方案中，C含量为0.19～0.25％。
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且强度较高的珠光体，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。
对于相变诱导塑性TRIP钢而言，目前的研究认为，锰元素在钢中起固溶强化和降低Ms点的作用，也可提高残余奥氏体的稳定性；也有研究认为，当钢中同时存在Si和Mn两种元素时，Si元素的存在会加剧Mn元素的偏聚程度，加强Mn对C原子的拖拽作用，推迟贝氏体的形成。锰含量较高时，会导致组织呈带状化、残余奥氏体过分稳定，不利于相变的发生；也会导致钢中晶粒粗化，增加钢的过热敏感性，当熔炼浇注与锻轧之后冷却不当时，容易使碳钢中产生白点。另外，增加Mn含量会增加合金成本，增加热处理前工序生产成本和生产难度。综合以上因素考虑，本发明将锰含量控制在1.6～2.4％范围之内。在一些实施方案中，Mn含量为1.6～2.2％。在一些实施方案中，Mn含量为1.8～2.4％。
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，硅可增加钢的冷加工变形硬化速率，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，而且在一定范围内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。但是钢中硅的含量过高，其塑性与韧性会显著下降。
对于相变诱导塑性TRIP钢而言，Si元素是铁素体形成元素，可以提高残余奥氏体的稳定性，也起到固溶强化的作用，提高钢的强度。硅元素还有缩小奥氏体相区，提高C元素在铁素体中活度的作用。较高的硅含量有利于获得较多的残余奥氏体，但过高的硅含量会使钢产生诸如坚硬的氧化层、差的表面性能、降低热轧钢板的润湿性、劣化表面质量等问题。硅元素的含量过低则不会带来稳定的令人满意的TRIP效应，所以硅含量必须控制在一定范围内。硅对奥氏体长大速率没有明显影响，但对形成的奥氏体形态和分布有明显影响。硅含量的增加将使得热处理前工序的制造难度增加，综合以上因素，本发明将硅含量控制在1.1～2.0％范围之内。在一些实施方案中，Si含量为1.1～1.7％。在一些实施方案中，Si含量为1.3～2.0％。
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，缩小奥氏体相区域，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。铬元素在提高钢的强度和硬度方面加入其他合金元素，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的含量≤0.3％。
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。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的含量≤0.3％。
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是强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。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。
对于TRIP钢而言，添加微合金化元素，能够通过晶粒细化和沉淀强化铁素体基体，还可以延迟贝氏体形成。贝氏体形成被延迟的原因是在冷却时强化铁素体的形成，这是显微组织晶粒细化后的结果。铁素体的形成导致残余奥氏体的碳富集，延迟了奥氏体转变为贝氏体，同时细小弥散的碳氮化物使贝氏体形核受到抑制，从而也延迟贝氏体形成。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：
第一，可在减少碳当量含量的同时提高强度，提高钢的焊接性能；
第二，将不纯物质如氧、氮、硫等固定起来，从而改善钢的可焊性；
第三，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，Nb的含量≤0.05％。在一些实施方案中，Ti的含量≤0.04％。在一些实施方案中，V的含量≤0.055％。
本发明通过快速热处理方法(包括快速加热、短时保温和快速冷却过程)精细化控制轧硬带钢在热处理过程中变形组织的回复、再结晶和相变过程，最终获得细小、均匀、弥散分布各项组织结构和良好的强塑性匹配。
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或长大不明显。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。
中国专利申请CN106811698B中公开的热处理工艺，其未对整个加热过程进行区分处理，且其加热过程采用的加热速率为20-60℃/s属于中等加热速率，是基于现有传统连续退火机组的加热技术实现的，无法根据材料组织转变的需要进行大范围的调控。
中国专利申请CN107794357B和美国专利申请US2019/0153558A1公开的热处理工艺中，虽然也对加热过程进行了分段处理：先以1-10℃/s的加热速率加热到300-500℃，然后以100-500℃/s的加热速率加热至单相奥氏体区850-950℃，保温不超过5s后水淬到室温。该处理方法要求必须将钢板加热到单相奥氏体的高温区，这提高了设备的耐高温要求，增加了制造难度，同时其采用水冷的冷却方式，虽然冷速极高，可大幅度减少晶粒组织在高温区间的长大时间，但是也不可避免的带来最终产品中合金元素分布不均匀，导致产品组织性能的不均匀和不稳定，冷速过高也会导致钢板板型不良和表面氧化等一系列问题。
只有通过综合控制整个热处理过程：包括快速加热(分区段控制加热速度)、短时均热和快速冷却过程，才能获得精细控制的最优的晶粒尺寸、合金元素和各相组织均匀分布，最终获得最优的强韧性匹配产品。
通过本发明的快速热处理方法后所获得的铁素体、奥氏体和贝氏体的多相组织，其平均晶粒尺寸在1～3μm，比现有传统技术生产的产品晶粒尺寸(通常在5～15μm)减小50～80％，通过晶粒细化可提高材料的强度，同时获得良好的塑性和韧性，提高材料的使用性能；而且本发明获得的铁素体、贝氏体和残余奥氏体组织主要为块状和颗粒状等多种形态，且分布更加均匀，从而变形阶段可获得更好的强塑性。
本发明所述的抗拉强度≥980MPa的低碳低合金TRIP钢的快速热处理制造方法，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
热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
3)冷轧
冷轧压下率40～80％，获得轧硬态带钢或钢板；
4)快速热处理
a)快速加热
冷轧后的带钢或钢板快速加热至770～860℃，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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；
b)均热
在奥氏体和铁素体两相区目标温度770～860℃进行均热，均热时间为40～120s；
c)冷却
将带钢或钢板均热后以5～15℃/s的冷却速率缓冷至670～770℃；随后以40～100℃/s(如50～100℃/s)的冷却速率快速冷却至贝氏体等温处理温度410～430℃；
d)贝氏体等温处理
在410～430℃温度区间进行贝氏体等温处理，等温处理时间150～250s；
e)等温处理结束后带钢或钢板以30～150℃/s的冷却速率冷却至室温。
优选的，所述快速热处理全过程用时为281～350s。
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，所述快速加热最终温度为790～830℃。
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30～300℃/s的加热速率从室温加热至550～625℃，第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃。
优选的，步骤4)中，所述快速加热采用两段式加热，第一段以50～300℃/s的加热速率从室温加热至550～625℃，第二段以80～300℃/s的加热速率从550～625℃加热至770～860℃。
优选的，步骤4)中，所述快速冷却过程冷却速率为50～100℃/s。
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860℃，降温后温度不低于770℃。
在一些实施方案中，本发明所述的抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的快速热处理制造方法，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
热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
3)冷轧
冷轧压下率40～80％，冷轧后获得轧硬态带钢或钢板；
4)快速热处理热镀锌
a)快速加热
将冷轧带钢或钢板由室温快速加热至770～860℃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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；
b)均热
在奥氏体和铁素体两相区目标温度770～860℃进行均热，均热时间为30～120s；
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40～100℃/s冷却速率快速冷却至410～430℃；
d)贝氏体等温处理
带钢或钢板在410～430℃进行贝氏体等温处理，等温处理时间60～150s；
e)再加热
等温处理结束后以10～30℃/s的加热速率加热至460～470℃；
f)热镀锌
随后将带钢或钢板浸入锌锅进行热镀锌；
g)带钢或钢板热镀锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀纯锌GI产品；或者，带钢或钢板热镀锌之后，以30～300℃/s如(30～100℃/s)的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～250℃/s(如30～100℃/s)的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
优选地，所述步骤4)的快速热处理全过程用时为118～328s，如118～238s。
优选地，步骤2)中，所述卷取温度为580～650℃。
优选地，步骤3)中，所述冷轧压下率为60～80％。
优选地，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。
优选地，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550～625℃，第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃。
优选地，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550～625℃，第二段以80～300℃/s的加热速率从550～625℃加热至770～860℃。
优选地，步骤4)中，所述快速加热最终温度为790～860℃。
优选地，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。
优选地，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860℃，降温后温度不低于770℃。
优选地，步骤4)中，所述带钢或钢板热镀锌之后，以30～200℃/s的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～200℃/s的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
在本发明所述的抗拉强度≥980MPa的低碳低合金TRIP钢的快速热处理制造方法中：
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 PCTCN2022084524-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。
传统热处理过程中受限于加热技术的影响均为慢速加热，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，且相变形核主要在已经长大的铁素体晶界处，形核率较低。因此最后得到的组织相对粗大。
快速加热条件下，变形基体刚刚完成再结晶或还没有完成再结晶(甚至还没有充分回复)，就开始发生铁素体向奥氏体的相转变，由于刚刚完成再结晶或还没有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。此时铁素体再结晶与奥氏体相变过程发生重叠，由于铁素体晶体内保留了大量位错等晶体缺陷，为奥氏体的形成提供了大量的形核点，使得奥氏体的形核呈现爆发式形核，因此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，因此奥氏体晶粒细小而且体积分数增大。
通过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50～500℃/s，采用两段式快速加热时加热速率定为15～500℃/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500～625℃，加热速率对回复过程的影响最大，控制加热速率为15～300℃/s，进一步优选为30～300℃/s；加热温度从500～625℃到奥氏体化温度770～860℃，加热速率对晶粒长大过程影响最大，控制加热速率为50～300℃/s；进一步优选为80～300℃/s。
2、均热温度控制
均热温度的选择需结合加热过程各温度阶段材料组织演变过程控制，同时需考虑后续快速冷却过程组织的演变和控制，这样才能最终获得优选的组织结构及分布。
均热温度通常取决于钢中C含量，传统的热处理工艺中一般将均热温度设置在A C1到A C3之间，或者A c3以上30～50℃。本发明利用快速加热技术在未充分回复及再结晶的基体中保留大量的位错等晶体缺陷，为奥氏体转变提供了充足的形核功，所以只需要将温度加热到A C1到A C3之间就可以获得较多的奥氏体。本发明中TRIP钢的C含量为：0.17～0.25％，A C1和A C3分别是730℃左右和870℃左右。TRIP钢中有大量未溶解的细小均匀分布的碳化物，在淬火加热过程中，能够对奥氏体颗粒的长大起到机械阻碍的作用，有利于细化合金钢的晶粒度，但是如果加热温度过高，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分布不均匀，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，合金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，以达到在冷却之后能够得到细小均匀的铁素体、贝氏体和残余奥氏体的目的。过高的均热温度会使奥氏体晶粒粗大，淬火过程中工件容易开裂，淬火后获得的组织也会较粗大，使钢的力学性能不佳。过低的均热温度，又会使奥氏体溶入的碳以及合金元素含量不足，令奥氏体碳浓度分布不均，使钢的淬透性大幅降低，对合金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770～860℃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，当钢中碳以及合金元素含量升高时，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。
由于本发明采用快速加热，在两相区材料含有大量位错，为奥氏体形成提供大量的形核点，并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，最后保留下来的残余奥氏体碳含量越多；但是如果保温时间过短，会使钢中合金元素分布不均，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其它合金元素的含量，当合金含量升高时，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。综上所述，本发明将均热保温时间定为40～120s。
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。
为了获得贝氏体，快冷时材料的冷速必须大于临界冷却速度才能够得到贝氏体，临界冷却速度主要取决于材料成分，本发明TRIP钢中的Si含量为1.1～2.0％，Mn含量为 1.6～2.4％，含量相对较高，所以Si和Mn很大程度加强了TRIP钢中奥氏体的淬透性，降低了临界冷却速度。
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。冷速太低则不能获得贝氏体组织，那么力学性能不能满足要求，但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)引起板形不良，甚至容易导致试样严重变形和开裂。所以本发明将快速冷却速度设置为40～100℃/s。
5、贝氏体等温处理温度
TRIP钢的贝氏体等温处理温度一般选择在贝氏体、铁素体与奥氏体自由能相等温度(T 0)以下，此时贝氏体自由能小于奥氏体自由能，自由能的降低为贝氏体转变提供了相变驱动力。由于钢铁材料化学组分不同，贝氏体等温处理温度也不相同，贝氏体等温温度一般选择在350～550℃之间。等温处理温度较高，原子扩散能力强，奥氏体部分转变成粒状贝氏体，并析出碳化物，降低过冷奥氏体的稳定性，残余奥氏体体积分数较低。较低温度等温处理时，需要原子扩散的贝氏体转变难以进行，可能导致发生无原子扩散的马氏体相变，马氏体为C的过饱和组织，在等温过程中，C扩散过于缓慢，难以在过冷奥氏体中富集，也会导致残余奥氏体体积分数减少，所以本发明贝氏体等温温度选择在410～430℃之间。
6、贝氏体等温处理时间
当贝氏体等温处理时间较短时贝氏体相变未能充分进行，C元素向奥氏体富集程度较低，奥氏体C含量较低其稳定性较差，在随后的冷却过程中，过冷奥氏体将转变为马氏体，马氏体组织具有高强度低延伸率的特点，因此对提高强塑性不利。随着等温时间的延长，贝氏体转变充分，本发明TRIP钢中贝氏体体积分数增加。等温时间延长，SEM显微组织变化不明显，贝氏体体积分数和形貌变化不大，此时主要为C元素向残余奥氏体富集的过程，随保温时间的延长，导致残余奥氏体碳含量升高，稳定性将增加，TRIP钢材料在服役使用过程中某形变区域的残余奥氏体随应变的发生可持续性地发生马氏体相变，该区域的材料强度增强，从而使应变可以转移到材料的其它区域，可显著提高材料延伸率。所以本发明将贝氏体等温时间设定得较长，在60～250s，如60～150s或150～250s。
7、热镀锌和合金化处理控制
对于高强度的热镀锌产品而言，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，有利于改善高强度热镀锌产品可镀性，减少表面漏镀缺陷，提高耐蚀性能，从而能提高成材率。
本发明通过对传统连续退火机组进行快速加热和快速冷却工艺改造，使其实现快速热处理工艺，可以极大的缩短退火炉加热及均热段的长度(较传统连续退火炉至少能缩短三分之一)，提高传统连续退火机组的生产效率，降低生产成本及能耗，显著减少连续退火炉炉辊数量，特别是减少高温炉段炉辊数量，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、材料过渡灵活、调控能力强等目的；对产品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制造难度，提高材料的成型、焊接等用户使用性能。
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，增加了再结晶晶粒和奥氏体晶粒的形核点，缩短晶粒长大时间，细化晶粒，所获得的TRIP钢的显微组织为贝氏体组织体积占比35～80％、铁素体组织体积占比15～60％、奥氏体组织体积占比5～18％的多相组织，且平均晶粒尺寸细化到1～3μm，比现有传统技术生产的产品晶粒尺寸(通常在5-15μm)减小50～80％；所获得的热镀锌TRIP钢的显微组织为贝氏体(35～80％)、铁素体(10～60％)和奥氏体(5～18％)的三相组织，平均晶粒尺寸为1～3μm，平均晶粒尺寸减小30～50％。所述钢中，贝氏体为亚微米级颗粒状；奥氏体为孤岛状分布的晶粒；贝氏体和奥氏体均匀分布在铁素体基体上；奥氏体还具有良好的热稳定性，-50℃奥氏体转化变率低于8％；-190℃奥氏体转变率低于30％，且不同大小、形状、取向的奥氏体可在不同应变条件下持续发生TRIP效应，显著提高材料性能。通过晶粒细化可提高材料的强度，同时获得良好的塑性和韧性，提高材料的使用性能。
(2)相比于传统热处理方式所得TRIP钢，该发明得到的TRIP钢为多相组织，且晶粒尺寸减小50～80％，材料的强韧性明显提高，其抗拉强度可控制在较小区间范围内(如1030～1120MPa，或1190～1300MPa)，产品力学性能的稳定性明显提高，且延伸率依然保持在较高水平(如21.3～24.5％)。相比于传统热处理方式所得相变诱导塑性TRIP钢，由于本发明得到的TRIP钢为晶粒细小的多相组织，平均晶粒尺寸为1～3μm，平均晶粒尺寸减小30～50％；可明显提高材料的强韧性，其屈服强度549～821MPa，抗拉强度提高至1030～1284MPa；延伸率18～24.4％；强塑积20.7～26.4GPa％。
(3)根据本发明所述的TRIP钢快速热处理工艺，热处理全过程用时可缩短至281～350s，根据本发明所述的热镀锌TRIP钢热处理工艺，热处理全过程最短用时可缩短至118s，大大降低了整个热处理工艺过程的时间(TRIP钢传统连续退火工艺时间通常在 9～11min)；特别是缩短了在600℃以上高温下的停留时间，从而提高了生产效率、减少了能耗，降低了生产成本。
(4)相比于传统的TRIP钢及其热处理工艺，本发明的快速热处理方法加热段和均热段时间缩短了60～80％，以及缩短了带钢在高温下的处理时间，整个热处理工序时间缩短至281～350s，可节能减排降耗，显著降低炉子设备的一次性投资，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)相比于传统的TRIP钢及其热处理工艺，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；另外由于产品晶粒细化和材料合金含量的减少，采用本发明技术得到的钢板扩孔性能和弯折性能等加工成形性能、焊接性能等用户使用性能也有所提高。
本发明得到的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP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及相应工业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所生产的TRIP钢显微组织图片。
图2是本发明实施例一试验钢A按传统工艺1所生产的TRIP钢显微组织图片。
图3是本发明实施例一试验钢P按实施例6所生产的TRIP钢显微组织图片。
图4是本发明实施例一试验钢M按实施例12所生产的TRIP钢显微组织图片。
图5是本发明实施例一试验钢G按实施例21所生产的TRIP钢显微组织图片。
图6是本发明实施例一试验钢S按实施例23所生产的TRIP钢显微组织图片。
图7是本发明实施例二试验钢A按实施例1所生产的TRIP钢显微组织图片。
图8是本发明实施例二试验钢A按传统工艺1所生产的TRIP钢显微组织图片。
图9是本发明实施例二试验钢P按实施例6所生产的TRIP钢显微组织图片。
图10是本发明实施例二试验钢M按实施例12所生产的TRIP钢显微组织图片。
图11是本发明实施例二试验钢G按实施例21所生产的TRIP钢显微组织图片。
图12是本发明实施例二试验钢S按实施例23所生产的TRIP钢显微组织图片。
图13是本发明实施例三试验钢A按实施例1所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图14是本发明实施例三试验钢A按传统工艺1下所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图15是本发明实施例三试验钢I按实施例17所生产的合金化热镀锌TRIP钢(GA)显微组织图片。
图16是本发明实施例三试验钢D按实施例22所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图17是本发明实施例三试验钢I按实施例34所生产的合金化热镀锌TRIP钢(GA)显微组织图片。
图18是本发明实施例四试验钢A按实施例1所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图19是本发明实施例四试验钢A按传统工艺1所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图20是本发明实施例四试验钢I按实施例17所生产的合金化热镀锌TRIP钢(GA)显微组织图片。
图21是本发明实施例四试验钢D按实施例22所生产的热镀纯锌TRIP钢(GI)显微组织图片。
图22是本发明实施例四试验钢I按实施例34所生产的合金化热镀锌TRIP钢(GA)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。
实施例一
本实施例试验钢的成分参见表1，本实施例及传统工艺的具体参数参见表2和表3，表4和表5为本发明试验钢成分按实施例及传统工艺制备所得钢的主要性能。
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良好匹配。通过本发明的方法获得的TRIP钢屈服 强度为549～716MPa，抗拉强度为1030～1120MPa，延伸率为21.3～24.5％，强塑积为23～25.9GPa％。
图1为典型成分A钢经过实施例1获得的组织图，图2为典型成分A钢经过传统工艺例1获得的组织图。从图上看，不同热处理方式处理后的组织存在非常大的区别。经过本发明实施例处理后A钢其组织(图1)主要为铁素体基体上弥散分布的均匀的贝氏体和残余奥氏体组织组成。通过本发明工艺处理后的组织：贝氏体、残余奥氏体组织都非常均匀的分布于铁素体基体中，这对提高材料强度和塑性都是非常有利的。而经过传统工艺处理的A钢组织(图2)则为典型的TRIP钢组织。大块铁素体晶界上分布着贝氏体和少量的残余奥氏体组织，平均晶粒尺寸大约在5-15μm左右。采用传统工艺处理的组织特点是：晶粒组织相对粗大，贝氏体及残余奥氏体组织含量相对较少，且分布相对不均匀。
图3为典型成分P钢经过实施例6获得的组织图，图4为典型成分M钢经过实施例12获得的组织图。图5为典型成分G钢经过实施例21获得的组织图，图6为典型成分S钢经过实施例23获得的组织图。实施例6、12、21、23为典型的快速热处理工艺。从图中可见，采用本发明方法，可以获得非常均匀、细小、弥散分布的铁素体、贝氏体和残余奥氏体各相组织。本发明的TRIP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。
Figure PCTCN2022084524-appb-000002
Figure PCTCN2022084524-appb-000003
Figure PCTCN2022084524-appb-000004
表4
Figure PCTCN2022084524-appb-000005
表5
Figure PCTCN2022084524-appb-000006
Figure PCTCN2022084524-appb-000007
实施例二
本实施例二试验钢的成分参见表6，本实施例及传统工艺的具体参数参见表7和表8，表9和表10为本发明试验钢成分按实施例及传统工艺制备所得钢的主要性能。
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良好匹配。通过本发明的方法获得TRIP钢的屈服强度为751～947MPa，抗拉强度为1193～1297MPa，延伸率为18.5～21％，强塑积为23～25.9GPa％。
图7为典型成分A钢经过实施例1获得的组织图，图8为典型成分A钢经过传统工艺例1获得的组织图。从图上看，不同热处理方式处理后的组织存在非常大的区别。经过本发明实施例处理后获得的TRIP钢其组织主要由铁素体、贝氏体、残余奥氏体及少量碳化物组成，且贝氏体、残余奥氏体组织都非常均匀的分布于铁素体基体中，这对提高材料强度和塑性都是非常有利的。而经过传统工艺处理获得的TRIP钢组织则为典型的TRIP钢组织，大块铁素体晶界上分布着贝氏体和少量的残余奥氏体组织，平均晶粒尺寸大约在5～15μm左右。采用传统工艺处理的组织特点是：晶粒组织相对粗大，贝氏体及残余奥氏体组织含量相对较少，且分布相对不均匀。
图9为典型成分P钢经过实施例6获得的组织图，图10为典型成分M钢经过实施例12获得的组织图。图11为典型成分G钢经过实施例21获得的组织图，图12为典型成分S钢经过实施例23获得的组织图。实施例6、12、21、23为典型的快速热处理工艺。从图中可见，采用本发明方法，可以获得非常均匀、细小、弥散分布的铁素体、贝氏体和残余奥氏体各相组织。本发明的TRIP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。
Figure PCTCN2022084524-appb-000008
Figure PCTCN2022084524-appb-000009
Figure PCTCN2022084524-appb-000010
表9
Figure PCTCN2022084524-appb-000011
表10
Figure PCTCN2022084524-appb-000012
从实施例一和二可知，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改造，使其实现快速热处理工艺，极大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，可以提高带钢表面质量的控制能力，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，使得连续热处理机组短小精悍、材料过渡灵活、而且调控能力强等优点；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制造难度，提高材料的焊接性能等用户使用性能。
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几十秒、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可以不超过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更柔性化的高强钢组织设计方法，进而在无需改变合金成分及轧制工艺等前工序条件前提下改善材料组织，提高材料性能。
以相变诱导塑性TRIP钢为代表的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应用价值，两者的结合必将会为TRIP钢的开发和生产提供更大的空间。
实施例三
本发明试验钢的成分参见表11，本实施例及传统工艺的具体参数参见表12(一段式加热)和表13(两段式加热)；表14和表15为本发明试验钢成分按表12和表13中实施例及传统工艺制备所得GI和GA热镀锌TRIP钢的主要性能。
从表11～表15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TRIP钢屈服强度可达549～620MPa，抗拉强度提高至1030～1164MPa；延伸率20.1～24.4％；强塑积20.7～25.8GPa％。
图13、图14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。从两张图上看，两种工艺热镀锌后的组织存在非常大的区别。经过本发明的快速热处理后的A钢其组织(图13)主要为细小铁素体基体上弥散分布的细小、均匀的贝氏体组织及碳化物组成，且贝氏体晶粒组织及碳化物都非常细小且均匀弥散分布，这对提高材料强度和塑性都是非常有利的。而经过传统工艺处理的A钢组织(图14)则为典型的TRIP钢组织图。即大块 白色铁素体组织晶界上存在少量的贝氏体和残余奥氏体组织。由于元素偏析等原因，传统工艺处理后的材料组织呈现一定的方向性，其组织沿轧制方向呈长条分布。采用传统工艺处理的组织特点是：晶粒大，且存在一定的带状组织，贝氏体及残余奥氏体沿铁素体晶界呈网状分布，铁素体晶粒相对粗大，铁素体及贝氏体两相组织分布不均匀。
图15为典型成分I钢经过实施例17(GA)获得的组织图，图16为典型成分D钢经过实施例22(GI)获得的组织图。图17为典型成分I钢经过实施例34(GA)获得的组织图。实施例17、22、34均为整个热处理周期较短的工艺。从图中可见，采用本发明快速热处理热镀锌方法，进行合金化处理后获得非常均匀、细小、弥散分布的各相组织，而通过传统工艺制备达到的钢带金相组织中，铁素体组织粗大，贝氏体及残余奥氏体组织分布在铁素体晶界上，为典型的热镀锌TRIP钢组织。因此本发明的热镀锌TRIP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。
Figure PCTCN2022084524-appb-000013
Figure PCTCN2022084524-appb-000014
Figure PCTCN2022084524-appb-000015
Figure PCTCN2022084524-appb-000016
表14
Figure PCTCN2022084524-appb-000017
表15
Figure PCTCN2022084524-appb-000018
实施例四
本实施例四试验钢的成分参见表16，本发明实施例及传统工艺的具体参数参见表17(一段式加热)和表18(两段式加热)；表19和表20为本发明试验钢成分按实施例及传统工艺制备所得GI热镀锌TRIP钢的主要性能，表19为本发明试验钢成分按表17和表18中实施例及传统工艺制备所得GI和GA热镀锌TRIP钢的主要性能。
从表16～表20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TRIP钢屈服强度可达771～821MPa，抗拉强度提高至1182～1284MPa；延伸率为18～22.2％；强塑积为22.6～26.4GPa％。
图18、图19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。从两张图上看，两种工艺热镀锌后的组织存在非常大的区别。经过本发明的快速热处理后的A钢其组织(图18)主要由细小、均匀切相互弥散分布的铁素体、贝氏体及残余奥氏体等相组织构成。各种组织非常细小且均匀弥散分布，这对提高材料强度和塑性都是非常有利的。
而经过传统工艺处理的A钢组织(图19)则为典型的TRIP钢组织图。采用传统工艺处理的组织特点是：组织沿轧向呈一定方向分布。铁素体组织占比稍大，贝氏体、残余奥氏体组织沿轧向分布，组织存在一定的不均匀性。
图20为典型成分I钢经过实施例17(GA)获得的组织图，图21为典型成分D钢经过实施例22(GI)获得的组织图。图22为典型成分I钢经过实施例34(GA)获得的组织图。实施例17、22、34均为整个热处理周期较短的工艺。从图中可见，采用本发明快速热处理热镀锌方法，进行合金化处理后获得非常均匀、细小、弥散分布的各相组织(图20)，而通过传统工艺9制备达到的钢带金相组织中，铁素体、贝氏体等组织均相对粗大，且具有一定的方向性。同时其组织中的铁素体组织含量相对较多，分布不均匀。因此本发明的热镀锌TRIP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。
Figure PCTCN2022084524-appb-000019
Figure PCTCN2022084524-appb-000020
Figure PCTCN2022084524-appb-000021
Figure PCTCN2022084524-appb-000022
Figure PCTCN2022084524-appb-000023
表19
Figure PCTCN2022084524-appb-000024
表20
Figure PCTCN2022084524-appb-000025
实施例三和四的结果显示，通过采用快速加热和快速冷却工艺对传统连续退火热镀锌机组进行工艺改造，使其实现快速热处理热镀锌工艺，可以极大的缩短传统连续退火热镀锌炉加热段及均热段的长度，提高传统连续退火热镀锌机组的生产效率，降低生产成本及能耗，减少连续退火热镀锌炉的炉辊数量，显著减少辊印、麻点、擦划伤等表面缺陷，因此提高了带钢表面质量的控制能力，容易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热镀锌工艺技术的新型连续退火机组，可实现热镀锌机组短小精悍、材料过渡灵活、调控能力强等优点；对热镀锌镀基板材料而言则可细化晶粒，进一步提高材料强度，降低合金成本及热处理前工序制造难度，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热镀锌工艺，对冷轧带钢的连续退火热镀锌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热镀锌炉子加热段长度，便于提高连续退火热镀锌机组的速度和生产效率，显著减少连续退火热镀锌机组炉内辊子数目，对于机组速度在180米/分左右的快速热处理热镀锌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热镀锌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。
以相变诱导塑性TRIP钢为代表的热镀锌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应用价值，两者的结合必将会为热镀锌TRIP钢的开发和生产提供更大的空间。

Claims (13)

  1. 抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其化学成分质量百分比为：C：0.17～0.25％，Si：1.1～2.0％，Mn：1.6～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；
    优选地，所述低碳低合金TRIP钢通过下述工艺获得：
    5)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    6)热轧、卷取
    热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
    7)冷轧
    冷轧压下率为40～80％；
    8)快速热处理
    冷轧后的钢板快速加热至770～860℃，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℃；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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；之后进行均热，均热温度为770～860℃，均热时间为40～120s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以40～100℃/s的冷却速率快速冷却至贝氏体等温处理温度410～430℃，并在此温度区间进行等温处理，等温处理时间150～250s；之后以30～150℃/s、优选30～100℃/s的冷却速率冷却至室温；
    优选地，所述低碳低合金热镀锌TRIP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
    热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
    3)冷轧
    冷轧压下率为40～80％；
    4)快速热处理热镀锌
    冷轧后的钢板快速加热至770～860℃，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550～625℃，第二段以50～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；之后进行均热，均热温度：770～860℃，均热时间：30～120s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至670～770℃，随后以40～100℃/s的冷却速率快速冷却至410～430℃；并在此温度区间进行等温处理，等温处理时间60～150s；等温处理结束后以10～30℃/s的加热速率加热至460～470℃，随后浸入锌锅进行热镀锌；
    热镀锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀纯锌GI产品；或者，热镀锌之后，以30～300℃/s(如30～100℃/s)的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～250℃/s(如30～100℃/s)的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
  2. 如权利要求1所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，所述的TRIP钢和热镀锌TRIP钢中，
    C的含量范围选自0.17～0.23％、0.19～0.21％、0.19～0.25％和0.21～0.23％；优选地，Si的含量范围选自：1.1～1.7％、1.3～1.5％、1.3～2.0％和1.5～1.8％；优选地，Mn的含量范围选自1.6～2.2％、1.8～2.0％、1.8～2.4％和2.0～2.2％；和/或
    Cr的含量≤0.3％，Mo的含量≤0.3％，Nb的含量≤0.05％，Ti的含量≤0.04％，V的含量≤0.055％。
  3. 如权利要求1或2所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，所述卷取温度为580～650℃，和/或，所述冷轧压下率为60～80％。
  4. 如权利要求1或2所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，
    所述低碳低合金TRIP钢的制备工艺中，所述快速热处理全过程用时为281～350s；所述低碳低合金热镀锌TRIP钢的制备工艺中，所述快速热处理和热镀锌全过程用时为118～328s，如118～238s；和/或
    所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    所述快速加热采用两段式加热，第一段以15～300℃/s或30～300℃/s的加热速率从室温加热至550～625℃；第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃；优选地，第一段以50～300℃/s的加热速率从室温加热至550～625℃；第二段以80～300℃/s的加热速率从550～625℃加热至770～860℃；和/或
    所述快速加热最终温度为790～860℃。
  5. 如权利要求1-4中任一项所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，所述带钢或钢板热镀锌之后，以30～200℃/s的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～200℃/s的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
  6. 如权利要求1-5中任一项所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，
    所述低碳低合金TRIP钢的金相组织为贝氏体、铁素体和奥氏体的多相组织，平均晶粒尺寸为1～3μm；优选地，贝氏体为亚微米级颗粒状，奥氏体为孤岛状分布的等轴晶粒，贝氏体和奥氏体均匀分布在铁素体基体上；优选地，所述多相组织中，贝氏体的体积比为35～80％、如35～45％或40～80％，铁素体的体积比10～60％、如10～50％，奥氏体的体积比为5～18％、如7～18％或5～15％；优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率为≥18％，强塑积≥23GPa％；优选地，该低碳低合金TRIP钢的奥氏体具有良好的热稳定性，-50℃奥氏体转化变率低于8％，-190℃奥氏体转变率低于30％；
    所述低碳低合金热镀锌TRIP钢的金相组织为贝氏体、铁素体和奥氏体的三相组织，平均晶粒尺寸为1～3μm；优选地，贝氏体为亚微米级颗粒状，奥氏体为孤岛状分布的等轴晶粒，贝氏体和奥氏体均匀分布在铁素体基体上；优选地，所述多相组织中，贝氏体的体积比为35～80％、如35～75％或40～80％，铁素体的体积比10～60％、如10～50％，奥氏体的体积比为5～18％、如7～18％或5～15％；优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率为≥18％，强塑积≥23GPa％；优选地，该低碳低合金TRIP钢的奥氏体具有良好的热稳定性，-50℃奥氏体转化变率低于8％，-190℃奥氏体转变率低于30％。
  7. 如权利要求1-6中任一项所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，
    所述抗拉强度≥980MPa的低碳低合金TRIP钢的化学成分质量百分比为：C：0.17～0.23％，Si：1.1～1.7％，Mn：1.6～2.2％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；优选地，C含量为0.19～0.21％；优选地，Si含量为1.3～1.5％；优选地，Mn含量为1.8～2.0％；优选地，所述低碳低合金TRIP钢的金相组织为贝氏体35～45％、铁素体10～60％和奥氏体5～15％的多相组织；优选地，该低碳低合金TRIP钢的屈服强度≥540MPa，抗拉强度≥980MPa，延伸率≥21％，强塑积≥23GPa％；更优选地，所述低碳 低合金TRIP钢的屈服强度为549～716MPa，抗拉强度为1030～1120MPa，延伸率为21.3～24.5％，强塑积23～25.9GPa％；或
    所述抗拉强度≥980MPa的低碳低合金TRIP钢化学成分质量百分比为：C：0.19～0.25％，Si：1.3～2.0％，Mn：1.8～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Ti、Nb、V中的一种或两种，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质；优选地，C含量为0.21～0.23％；优选地，Si含量为1.5～1.8％；优选地，Mn含量为2.0～2.2％；优选地，该低碳低合金TRIP钢的金相组织为贝氏体40～80％、铁素体10～50％、奥氏体7～18％的三相组织；优选地，该低碳低合金TRIP钢的屈服强度≥770MPa，抗拉强度≥1180MPa，延伸率≥17％，强塑积≥23GPa％；更优选地，该低碳低合金TRIP钢的为776～861MPa，抗拉强度为1180～1297MPa，延伸率为17～21％，强塑积为23～25.9GPa％。
  8. 如权利要求1-6中任一项所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢，其特征在于，
    所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的化学成分质量百分比为：C：0.17～0.23％，Si：1.1～1.7％，Mn：1.6～2.2％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，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，C含量为0.19～0.21％；优选地，Si含量为1.3～1.5％；优选地，Mn含量为1.8～2.0％；优选地，所述快速热处理、热镀锌全过程用时为118～328s；优选地，所述低碳低合金热镀锌TRIP钢的金相组织为贝氏体35～75％、铁素体10～60％、奥氏体5～15％的三相组织；优选地，所述低碳低合金热镀锌TRIP钢的屈服强度549～620MPa，抗拉强度提高至1030～1164MPa；延伸率20.1～24.4％；强塑积20.7～25.8GPa％；或
    所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的化学成分质量百分比为：C：0.19～0.25％，Si：1.3～2.0％，Mn：1.8～2.4％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，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，C含量为0.21～0.23％；优选地，Si含量为1.5～1.9％；优选地，Mn含量为2.0～2.2％；优选地，所述快速热处理、热镀锌全过程用时为118～238s；优选地，所述低碳低合金热镀锌TRIP钢的金相组织为贝氏体40～80％、铁素体10～50％、奥氏体7～18％的三相组织；优选地，所述低碳低合金热镀锌TRIP钢的屈服强度为771～821MPa，抗拉强度提高为1182～1284MPa；延伸率为18～22.2％；强塑积为22.6～26.4GPa％。
  9. 权利要求1-8中任一项所述的抗拉强度≥980MPa的低碳低合金TRIP钢或抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的制造方法，其特征在于，
    所述抗拉强度≥980MPa的低碳低合金TRIP钢的制造方法包括以下步骤：
    5)冶炼、铸造
    按上述化学成分冶炼并铸造成板坯；
    6)热轧、卷取
    热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
    7)冷轧
    冷轧压下率40～80％，获得轧硬态带钢或钢板；
    8)快速热处理
    f)快速加热
    冷轧后的带钢或钢板快速加热至770～860℃，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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；
    g)均热
    在奥氏体和铁素体两相区目标温度770～860℃进行均热，均热时间为40～120s；
    h)冷却
    将带钢或钢板均热后以5～15℃/s的冷却速率缓冷至670～770℃；随后以40～100℃/s(如50～100℃/s)的冷却速率快速冷却至贝氏体等温处理温度410～430℃；
    i)贝氏体等温处理
    在410～430℃温度区间进行贝氏体等温处理，等温处理时间150～250s；
    j)等温处理结束后带钢或钢板以30～150℃/s的冷却速率冷却至室温；
    所述抗拉强度≥980MPa的低碳低合金热镀锌TRIP钢的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
    热轧终轧温度≥A r3，随后冷却至550～680℃进行卷取；
    3)冷轧
    冷轧压下率40～80％，冷轧后获得轧硬态带钢或钢板；
    4)快速热处理热镀锌
    a)快速加热
    将冷轧带钢或钢板由室温快速加热至770～860℃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550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～860℃；
    b)均热
    在奥氏体和铁素体两相区目标温度770～860℃进行均热，均热时间为30～120s；
    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40～100℃/s冷却速率快速冷却至410～430℃；
    d)贝氏体等温处理
    带钢或钢板在410～430℃进行贝氏体等温处理，等温处理时间60～150s；
    e)再加热
    等温处理结束后以10～30℃/s的加热速率加热至460～470℃；
    f)热镀锌
    随后将带钢或钢板浸入锌锅进行热镀锌；
    g)带钢或钢板热镀锌之后，以30～150℃/s的冷却速率快速冷却至室温，获得热镀纯锌GI产品；或者，带钢或钢板热镀锌之后，以30～300℃/s如(30～100℃/s)的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～250℃/s(如30～100℃/s)的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
  10. 如权利要求9所述的方法，其特征在于，所述卷取温度为580～650℃。
  11. 如权利要求9或10所述的方法，其特征在于，所述冷轧压下率为60～80％。
  12. 如权利要求9～11中任一项所述的方法，其特征在于，
    所述低碳低合金TRIP钢的制备工艺中，所述快速热处理全过程用时为281～350s；所述低碳低合金热镀锌TRIP钢的制备工艺中，所述快速热处理和热镀锌全过程用时为118～328s，如118～238s；和/或
    所述快速加热采用一段式加热时加热速率为50～300℃/s；和/或
    所述快速加热采用两段式加热，第一段以15～300℃/s或30～300℃/s的加热速率从室温加热至550～625℃；第二段以50～300℃/s的加热速率从550～625℃加热至770～860℃；优选地，第一段以50～300℃/s的加热速率从室温加热至550～625℃；第二段以80～300℃/s 的加热速率从550～625℃加热至770～860℃；和/或
    所述快速加热最终温度为790～860℃。
  13. 如权利要求9～12中任一项所述的方法，其特征在于，所述带钢或钢板热镀锌之后，以30～200℃/s的加热速率加热到480～550℃进行合金化处理，合金化处理时间5～20s；合金化处理后以30～200℃/s的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
PCT/CN2022/084524 2021-04-02 2022-03-31 抗拉强度≥980MPa的低碳低合金TRIP钢或热镀锌TRIP钢及其制造方法 WO2022206912A1 (zh)

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