WO2022206911A1 - 抗拉强度≥1180MPa的低碳低合金Q＆P钢或热镀锌Q＆P钢及其制造方法 - Google Patents

抗拉强度≥1180MPa的低碳低合金Q＆P钢或热镀锌Q＆P钢及其制造方法 Download PDF

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WO2022206911A1
WO2022206911A1 PCT/CN2022/084518 CN2022084518W WO2022206911A1 WO 2022206911 A1 WO2022206911 A1 WO 2022206911A1 CN 2022084518 W CN2022084518 W CN 2022084518W WO 2022206911 A1 WO2022206911 A1 WO 2022206911A1
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steel
low
hot
carbon
alloy
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French (fr)
Inventor
李俊
王健
刘赓
王骏飞
毛展宏
杜小峰
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Baoshan Iron and Steel Co Ltd
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Baoshan Iron and Steel Co Ltd
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Priority claimed from CN202110360562.7A external-priority patent/CN115181898B/zh
Priority claimed from CN202110360154.1A external-priority patent/CN115181887B/zh
Priority claimed from CN202110360131.0A external-priority patent/CN115181884B/zh
Priority claimed from CN202110360528.XA external-priority patent/CN115181895B/zh
Application filed by Baoshan Iron and Steel Co Ltd filed Critical Baoshan Iron and Steel Co Ltd
Priority to US18/551,266 priority Critical patent/US20240167130A1/en
Priority to JP2023560448A priority patent/JP7734205B2/ja
Priority to EP22779091.2A priority patent/EP4317511A4/en
Priority to KR1020237032119A priority patent/KR20230166081A/ko
Publication of WO2022206911A1 publication Critical patent/WO2022206911A1/zh
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    • 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
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    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention belongs to the technical field of rapid heat treatment of materials, and particularly relates to a low-carbon and low-alloy Q&P steel or a low-carbon and low-alloy hot-dip galvanized Q&P steel with a tensile strength of ⁇ 1180 MPa and a manufacturing method thereof.
  • the Q&P heat treatment process is a new type of continuous heat treatment technology proposed by Speer et al. in the early 21st century.
  • the process mainly includes four steps:
  • the sample is rapidly cooled to a temperature between Ms s and M f to obtain a dual-phase structure mainly composed of martensite and retained austenite;
  • Q&P steel is essentially a martensitic steel, but it is different from traditional tempered martensitic steel. Under the same strength as tempered martensitic steel, the plasticity of Q&P steel is greatly improved. This is due to the existence of retained austenite in the structure of Q&P steel, which is transformed into martensite during deformation, resulting in the so-called TRIP effect, which greatly improves the plasticity of the steel.
  • the US patent application US2003/027825 proposes the general process of the Q&P steel production process, and limits the austenitization process to be carried out at high temperatures, and the material structure needs to be fully austenitized. For the actual production process, this temperature is too high (850 -950°C), and the time is long (usually, the austenitization process of the steel plate needs to be kept for 2 to 5 minutes), the equipment requirements are higher, and the manufacturing cost is also higher.
  • Chinese patent CN1081931138B discloses "a 980MPa grade automotive cold-rolled high-strength Q&P steel and its production method", the chemical composition of the steel is: C: 0.18-0.24%, Si: 0.6-1.3%, Mn: 1.6-2.4 %, P: 0.02 ⁇ 0.04%, S ⁇ 0.005%, Nb: 0.04 ⁇ 0.07%, N ⁇ 0.006%, Als: 0.5 ⁇ 1.0%, the balance is Fe and other unavoidable impurity elements.
  • the final rolling temperature of the hot rolling process is 870-910°C, and the coiling temperature is 660-710°C; the cold-rolling reduction ratio of the cold-rolling process is ⁇ 45%; the soaking temperature of the continuous annealing process is 770-840°C , the holding temperature of the over-aging section is 300-440°C, the holding time of the soaking section is 60-225 s, and the holding time of the over-aging section is 300-1225 s;
  • the yield strength of the obtained steel plate is greater than 550MPa, the tensile strength is greater than 980MPa, and the elongation after fracture is greater than 18%.
  • the main feature of this invention steel is to obtain the result of good strong-plasticity coordination through the traditional Q&P process. Due to the traditional heat treatment process, the soaking time and the distribution time are long, and the alloy content is relatively high, which will also increase the manufacturing cost and reduce the manufacturing flexibility.
  • Chinese patent application CN109136779A discloses "a preparation method of martensitic matrix 1100MPa grade rare earth Q&P steel", the chemical composition mass percentage of the invention steel is: C: 0.15-0.22%, Si: 0.6-1.7%, Mn: 1.1-2.4 %, Mo: 0.1-0.5%, Al: 0.1-0.5%, V: 0.05-0.11%, Y: 0.01-0.05%, P: 0.02-0.04%, S ⁇ 0.005%, Nb: 0.04-0.07%, N ⁇ 0.006%, B: 0.001-0.006%, the balance is Fe and other unavoidable impurity elements.
  • the tensile strength of the obtained steel sheet is about 1100 MPa, and the elongation after fracture is about 20%.
  • the main feature of the steel of the invention is that the rare earth Y and alloy elements such as Mo, V, Nb are added simultaneously to refine the grains, and the content of the Mn element is reduced to improve the welding performance.
  • the manufacturing process requires two castings. Smelting process: according to the composition formula given by the invention, after batching, it undergoes converter smelting, secondary refining and casting in a vacuum furnace to obtain a casting billet; trace element melting process: adding trace alloy element powder (Mo, Al) in the electric arc melting furnace. , V, Y, Nb, N, B, etc.) to obtain a secondary slab.
  • Hot rolling process use a heating furnace to heat the billet to 1100-1150 °C and keep it for 1-3 hours, and then perform hot rolling.
  • the final rolling temperature is 820-880 °C, and the coiling temperature is 550-650 °C.
  • the thickness of the obtained steel plate It is 1.5-3.0mm, and then water quenched to room temperature; cold rolling process: after pickling, multi-pass cold rolling is performed to obtain a steel plate with a thickness of 1.2-1.5mm;
  • the third time secondary carbon distribution process: the material is subjected to secondary carbon distribution for 10-300s at a certain temperature T 1 between MS and M f (the temperature of T 1 is slightly lower than T 0 ), and then the material is water quenched to room temperature.
  • the invention has complex manufacturing process, high energy consumption, high and complex alloy content, and multiple water quenching treatments, which involve the removal of the oxide layer on the surface of the material, which brings about many problems in terms of environment and energy consumption, resulting in increased manufacturing costs and increased manufacturing costs. Reduced flexibility.
  • Chinese patent application CN108431248A discloses "a method for manufacturing a high-strength steel sheet with improved ductility and formability and the obtained steel sheet", the chemical composition mass percentage of the invention steel is: C: 0.15-0.23%, Mn : 2.0 ⁇ 2.8%, Si: 1.0 ⁇ 2.1%, Al: 0.02 ⁇ 1.0%, Al+Si: 1.0 ⁇ 2.1%, Nb: 0 ⁇ 0.035%, Mo ⁇ 0.3%, Cr ⁇ 0.04%, the balance is Fe and other unavoidable impurity elements.
  • the steel sheet is annealed at the annealing temperature TA to obtain a structure comprising at least 65% austenite and up to 35% ferrite.
  • the tensile strength of the obtained steel plate is greater than 1180MPa, and the elongation after fracture is greater than 12%.
  • the main feature of the invention steel is that high Mn, high Si and high Al components are used to control the proportion of each phase in the final structure through traditional Q&P process, so as to obtain the result of good strong-plastic coordination. Due to the traditional heat treatment process, the soaking time and the dispensing time are long, which increases the manufacturing cost and reduces the manufacturing flexibility.
  • 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 manufacturing steps of this invention steel are as follows:
  • the temperature range for reheating of the forging billet in step 1) is 1100-1200°C, the holding time is 3-5h, the rolling temperature of hot rolling is 1050-1150°C, and the final rolling temperature is 850-900°C;
  • the 4-high reversible rolling mill performs 7 passes of reciprocating rolling, the reduction rate of the first two passes is 30-50%, and the reduction rate of the last five passes is 20-30%, and then the water is cooled to 650-750 °C and then put into The asbestos is kept for 8-10h to simulate the curling process, and the thickness of the hot-rolled strip is 4-5.5mm.
  • the cold rolling described in step 2) is unidirectionally rolled using a four-high rolling mill, and the rolling passes are 10-15 passes, including 3-5 passes of temper rolling, and the final cold-rolled strip thickness is 1.0- 1.5mm.
  • the austenitization temperature of the cold-rolled strip steel in step 3) is 870-920° C., and the austenitization holding time is 5-15 minutes.
  • Step 4) The removal thickness of the iron oxide scale and the decarburized layer is 50-100 ⁇ m on the upper and lower bottom surfaces, the annealing temperature for reheating the pre-quenched strip steel is 780-830° C., and the annealing holding time is 3-8 min. Then carry out salt bath cooling, the salt bath cooling speed is 100-200°C/s, the salt bath heat preservation temperature is 320-400°C, and the heat preservation time is 5-10min.
  • the main feature of the invention steel is that the grains are refined by adding more microalloying elements Nb and Ti to obtain high elongation (A% ⁇ 24%) and high strength ( ⁇ 980MPa).
  • the invention adopts the method of performing two heat treatments on the cold-rolled steel strip: the cold-rolled steel strip after pickling and cold-rolling treatment is firstly subjected to a complete austenitizing annealing, and then quenched into a complete annealing process. Martensitic structure, followed by surface dephosphorization and decarburization layer removal, and then reheating and annealing, and finally the finished strip steel is obtained.
  • the method has problems such as high addition of microalloying elements and increased manufacturing cost and increased difficulty of manufacturing procedures caused by two annealing times.
  • Chinese patent 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 material has high strength and large deformation resistance, which brings great difficulties for subsequent pickling and cold rolling production;
  • the third is the soaking time of 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. .
  • 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 410-430°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;
  • this 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 material structure of different temperature sections, but all of them are processed at a temperature of 40-300°C/s. 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 weaken the grain refinement effect produced by rapid heating, which is very unfavorable for the improvement of material strength and toughness. ;
  • 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 it does not play its due role, so it is unnecessary.
  • 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 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 dual-phase steel is calculated by weight percentage: 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 are 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 reheated to 850-950 °C in the single-phase austenite region; after that, the steel plate is water-cooled to room temperature immediately after the heat preservation does not exceed 5s 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 more alloying elements, and the yield strength and tensile strength both exceed 1000MPa, so this process is not suitable for heat treatment.
  • the manufacturing process before heat treatment and subsequent user use bring great 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 no high practical application value or little practical application value.
  • Chinese patent CN1081931138B discloses "a 980MPa grade automotive cold-rolled high-strength Q&P steel and its production method", the chemical composition of the steel is: C: 0.18-0.24%, Si: 0.6-1.3%, Mn: 1.6-2.4 %, P: 0.02 ⁇ 0.04%, S ⁇ 0.005%, Nb: 0.04 ⁇ 0.07%, N ⁇ 0.006%, Als: 0.5 ⁇ 1.0%, the balance is Fe and other unavoidable impurity elements.
  • the final rolling temperature of the hot rolling process is 870-910°C, and the coiling temperature is 660-710°C; the cold-rolling reduction ratio of the cold-rolling process is ⁇ 45%; the soaking temperature of the continuous annealing process is 770-840°C , the holding temperature of the over-aging section is 300-440°C, the holding time of the soaking section is 60-225 s, and the holding time of the over-aging section is 300-1225 s;
  • the yield strength of the obtained steel plate is greater than 550MPa, the tensile strength is greater than 980MPa, and the elongation after fracture is greater than 18%.
  • the main feature of this invention steel is to obtain the result of good strong-plasticity coordination through the traditional Q&P process. Due to the traditional heat treatment process, the soaking time and the distribution time are long, and the alloy content is relatively high, which will also increase the manufacturing cost and reduce the manufacturing flexibility.
  • Chinese patent application CN109136779A discloses "a preparation method of martensitic matrix 1100MPa grade rare earth Q&P steel", the chemical composition mass percentage of the invention steel is: C: 0.15-0.22%, Si: 0.6-1.7%, Mn: 1.1-2.4 %, Mo: 0.1-0.5%, Al: 0.1-0.5%, V: 0.05-0.11%, Y: 0.01-0.05%, P: 0.02-0.04%, S ⁇ 0.005%, Nb: 0.04-0.07%, N ⁇ 0.006%, B: 0.001-0.006%, the balance is Fe and other unavoidable impurity elements.
  • the tensile strength of the obtained steel sheet is about 1100 MPa, and the elongation after fracture is about 20%.
  • the main feature of the steel of the invention is that the rare earth Y and alloy elements such as Mo, V, Nb are added simultaneously to refine the grains, and the content of the Mn element is reduced to improve the welding performance.
  • the manufacturing process requires two castings. Smelting process: according to the composition formula given by the invention, after batching, it undergoes converter smelting, secondary refining and casting in a vacuum furnace to obtain a casting billet; trace element melting process: adding trace alloy element powder (Mo, Al) in the electric arc melting furnace. , V, Y, Nb, N, B, etc.) to obtain a secondary slab.
  • Hot rolling process use a heating furnace to heat the billet to 1100-1150 °C and keep it for 1-3 hours, and then perform hot rolling.
  • the final rolling temperature is 820-880 °C, and the coiling temperature is 550-650 °C.
  • the thickness of the obtained steel plate It is 1.5-3.0mm, and then water quenched to room temperature; cold rolling process: after pickling, multi-pass cold rolling is performed to obtain a steel plate with a thickness of 1.2-1.5mm;
  • the third time secondary carbon distribution process: the material is subjected to secondary carbon distribution for 10-300 s at a certain temperature T1 between MS and Mf (T1 temperature is slightly lower than T0), and then the material is water quenched to room temperature.
  • the invention has complex manufacturing process, high energy consumption, high and complex alloy content, and multiple water quenching treatments, which involve the removal of the oxide layer on the surface of the material, which brings about many problems in terms of environment and energy consumption, resulting in increased manufacturing costs and increased manufacturing costs. Reduced flexibility.
  • Chinese patent application CN108431248A discloses "a method for manufacturing a high-strength steel sheet with improved ductility and formability and the obtained steel sheet", the chemical composition mass percentage of the invention steel is: C: 0.15-0.23%, Mn : 2.0 ⁇ 2.8%, Si: 1.0 ⁇ 2.1%, Al: 0.02 ⁇ 1.0%, Al+Si: 1.0 ⁇ 2.1%, Nb: 0 ⁇ 0.035%, Mo ⁇ 0.3%, Cr ⁇ 0.04%, the balance is Fe and other unavoidable impurity elements.
  • the steel sheet is annealed at the annealing temperature TA to obtain a structure comprising at least 65% austenite and up to 35% ferrite.
  • the tensile strength of the obtained steel plate is greater than 1180MPa, and the elongation after fracture is greater than 12%.
  • the main feature of the invention steel is that high Mn, high Si and high Al components are used to control the proportion of each phase in the final structure through traditional Q&P process, so as to obtain the result of good strong-plastic coordination. Due to the traditional heat treatment process, the soaking time and the dispensing time are long, which increases the manufacturing cost and reduces the manufacturing flexibility.
  • 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 manufacturing steps of this invention steel are as follows:
  • the temperature range for reheating of the forging billet in step 1) is 1100-1200°C, the holding time is 3-5h, the rolling temperature of hot rolling is 1050-1150°C, and the final rolling temperature is 850-900°C;
  • the 4-high reversible rolling mill performs 7 passes of reciprocating rolling, the reduction rate of the first two passes is 30-50%, and the reduction rate of the last five passes is 20-30%, and then the water is cooled to 650-750 °C and then put into The asbestos is kept for 8-10h to simulate the curling process, and the thickness of the hot-rolled strip is 4-5.5mm.
  • the cold rolling described in step 2) is unidirectionally rolled using a four-high rolling mill, and the rolling passes are 10-15 passes, including 3-5 passes of temper rolling, and the final cold-rolled strip thickness is 1.0- 1.5mm.
  • the austenitization temperature of the cold-rolled strip steel in step 3) is 870-920° C., and the austenitization holding time is 5-15 minutes.
  • Step 4) The removal thickness of the iron oxide scale and the decarburized layer is 50-100 ⁇ m on the upper and lower bottom surfaces, the annealing temperature for reheating the pre-quenched strip steel is 780-830° C., and the annealing holding time is 3-8 min. Then carry out salt bath cooling, the salt bath cooling speed is 100-200°C/s, the salt bath heat preservation temperature is 320-400°C, and the heat preservation time is 5-10min.
  • the main feature of the invention steel is that the grains are refined by adding more microalloying elements Nb and Ti to obtain high elongation (A% ⁇ 24%) and high strength ( ⁇ 980MPa).
  • the invention adopts the method of performing two heat treatments on the cold-rolled steel strip: the cold-rolled steel strip after pickling and cold-rolling treatment is firstly subjected to a complete austenitizing annealing, and then quenched into a complete annealing process. Martensitic structure, followed by surface dephosphorization and decarburization layer removal, and then reheating and annealing, and finally the finished strip steel is obtained.
  • the method has problems such as high addition of microalloying elements and increased manufacturing cost and increased difficulty of manufacturing procedures caused by two annealing times.
  • the related researches on cold-rolled Q&P steel products and annealing process are based on the heating rate (5 ⁇ 20°C/s) of the existing industrial equipment to slowly heat the strip to make it The recrystallization and austenitizing transformation are completed in sequence, so the heating and soaking time is relatively long and the energy consumption is high.
  • the traditional continuous annealing production line also has a large number of rolls in the high-temperature furnace section of the strip, etc.
  • the traditional continuous annealing unit According to the product outline and production capacity requirements, the general soaking time is required to be 1 to 3 minutes. For a traditional production line with a unit speed of about 180 m/min, the number of rollers in the high temperature furnace section is generally 20 to 40, so that the belt Steel surface quality control is more difficult.
  • the purpose of the present invention is to provide a low-carbon and low-alloy Q&P steel with a tensile strength of ⁇ 1180 MPa, a low-carbon and low-alloy hot-dip galvanized Q&P steel with a tensile strength of ⁇ 1180 MPa, and a method for producing them by rapid heat treatment.
  • the invention changes the recovery, recrystallization and austenite transformation process of the deformed structure through rapid heat treatment, increases the nucleation rate (including the recrystallization nucleation rate and the austenite phase deformation nucleation rate), shortens the grain growth time, and reduces the size of the grains. Grain, increase the retained austenite content, thereby further improving the strength and plasticity of the material.
  • the matrix structure of the low-carbon and low-alloy Q&P steel of the present invention is evenly distributed, obvious lamellar tempered martensite appears, the grain size is 1-3 ⁇ m, and there are evenly distributed retained austenite around the martensite-strengthening phase grains.
  • Phase and ferrite phase in which lamellar structure accounts for 75-90% of martensite structure, 10-25% of retained austenite structure, and 3-10% of ferrite structure according to volume fraction.
  • the low-carbon and low-alloy Q&P steel of the present invention has a yield strength of ⁇ 660 MPa, a tensile strength of ⁇ 1180 MPa, an elongation of ⁇ 18%, and a strong-plastic product of ⁇ 24 GPa%, and has good strength and toughness matching and user performance such as forming and welding.
  • the invention adopts the rapid heat treatment process to improve the production efficiency and reduce the alloy content in the steel of the same grade, thereby reducing the production cost and the manufacturing difficulty of the pre-heat treatment process, significantly reducing the number of furnace rolls and improving the surface quality of the material.
  • the technical scheme of the present invention is:
  • Low-carbon low-alloy Q&P steel with tensile strength ⁇ 1180MPa its chemical composition mass percentage is: C: 0.16-0.23%, Si: 1.1-2.0%, Mn: 1.6-3.0%, P ⁇ 0.015%, S ⁇ 0.005% , Al: 0.02 to 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 of impurities.
  • the metallographic structure of the low-carbon low-alloy Q&P steel with tensile strength ⁇ 1180MPa is a multiphase structure of 75-90% martensite, 10-25% retained austenite, and 3-10% ferrite,
  • the matrix structure is evenly distributed, with obvious lamellar tempered martensite, and the grain size is 1-3 ⁇ m.
  • the austenite in the metallographic structure of the Q&P steel has good thermal stability, the -50°C austenite transformation rate is lower than 8%, and the -190°C austenite transformation rate is lower than 30%.
  • the Q&P steel has a yield strength of 668-1112 MPa, a tensile strength of 1181-1350 MPa, an elongation of 18.9-24.2%, and a strong-plastic product of 24.1-28.97 GPa%.
  • the content of C is selected from the range of 0.17-0.23%, 0.19-0.21% and 0.18-0.21%.
  • the content of Si is selected from the range of 1.1-1.7%, 1.3-1.5%, 1.4-2.0% and 1.6-1.8%.
  • the content of Mn is selected from the range of 1.6-2.2%, 1.8-2.0%, 2.4-3.0% and 2.6-2.8%.
  • the content of Cr is ⁇ 0.35%, such as ⁇ 0.25%; the content of Mo is ⁇ 0.25%; the content of Nb is ⁇ 0.06%, such as ⁇ 0.04%; The content of Ti is ⁇ 0.065%, such as ⁇ 0.04%, such as 0.006-0.016%; the content of V is ⁇ 0.055%, such as ⁇ 0.035%.
  • the low-carbon and low-alloy Q&P steel with a tensile strength of ⁇ 1180 MPa according to the present invention is obtained by the following process:
  • the cold rolling reduction rate is 40-85%;
  • the cold-rolled steel plate is rapidly heated to 770-845°C, and the rapid heating adopts one-stage or two-stage type; when one-stage rapid heating is used, the heating rate is 50-500°C/s; when two-stage rapid heating is used, the first 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 from 550-625°C to 770-845°C at a heating rate of 50-500°C/s; then soaking is performed , the soaking temperature is 770 ⁇ 845°C, and the soaking time is 10 ⁇ 60s;
  • the hot rolling finishing temperature is greater than or equal to Ar3.
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the entire process of the rapid heat treatment in step 4) takes 71 to 186 s.
  • 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; 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 845 °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; 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 845 °C.
  • the rapid cooling rate of the steel plate is 50-150°C/s.
  • the chemical composition mass percentage of the low-carbon low-alloy Q&P steel with tensile strength ⁇ 1180MPa of the present invention is: C: 0.17-0.23%, Si: 1.1-1.7%, Mn: 1.6-2.2%, P ⁇ 0.015%, S ⁇ 0.005%, 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 metallographic structure of the Q&P steel is a multiphase structure of 75-85% martensite, 10-25% retained austenite and 3-10% ferrite, the matrix structure is uniformly distributed, and obvious lamellae appear.
  • Tempered martensite the grain size is 1-3 ⁇ m, there is a uniformly distributed ferrite phase around the grains of the martensitic strengthening phase, and the grains of the martensite strengthening phase are mainly flaky.
  • the austenite in the metallographic structure of the Q&P steel has good thermal stability, the -50°C austenite transformation rate is lower than 8%, and the -190°C austenite transformation rate is lower than 30%.
  • the Q&P steel has a yield strength of 668-1002 MPa, a tensile strength of 1181-1296 MPa, an elongation of 18.9-24.2%, and a strong-plastic product of 24.1-28.6 GPa%.
  • the mass percentage of chemical composition of the low-carbon low-alloy Q&P steel with tensile strength ⁇ 1180MPa of the present invention is: C: 0.16-0.23%, Si: 1.4-2.0%, Mn: 2.4-3.0%, Ti : 0.006 ⁇ 0.016%, P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, can also contain one or two of Cr, Mo, Nb, V, and, Cr+Mo+Ti+Nb +V ⁇ 0.5%, the balance is Fe and other inevitable impurities.
  • the low-carbon and low-alloy Q&P steel with a tensile strength of ⁇ 1180 MPa is a low-carbon and low-alloy Q&P steel with a tensile strength of ⁇ 1280 MPa.
  • the C content is 0.18-0.21%.
  • the Si content is 1.6-1.8%.
  • the Mn content is 2.6-2.8%.
  • the metallographic structure of this Q&P steel is a multiphase structure of 80-90% martensite, 10-20% retained austenite, and 3-5% ferrite.
  • the matrix structure is uniformly distributed, and there is obvious lamellar tempering. Martensite, the grain size is 1-3 ⁇ m, there is a uniform distribution of ferrite phase around the grains of the martensite strengthening phase, and the grains of the martensite strengthening phase are mainly flaky structure.
  • the austenite in the metallographic structure of the Q&P steel has good thermal stability, the -50°C austenite transformation rate is lower than 8%, and the -190°C austenite transformation rate is lower than 30%.
  • the Q&P steel has a yield strength of 754-1112 MPa, a tensile strength of 1281-1350 MPa, an elongation of 19-22.2%, and a strength-plastic product of 24.8-28.97 GPa%.
  • Another aspect of the present invention provides a low-carbon and low-alloy hot-dip galvanized Q&P steel with a tensile strength of ⁇ 1180 MPa, the chemical composition mass percentages of which are: C: 0.16-0.23%, Si: 1.1-2.0%, Mn: 1.6-3.0%, P ⁇ 0.015%, S ⁇ 0.005%, preferably ⁇ 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 metallographic structure of the hot-dip galvanized Q&P steel is a three-phase structure of martensite, ferrite and austenite, the matrix structure is uniformly distributed, lamellar tempered martensite appears, and the grain size is It is 1-3 ⁇ m, and there is a uniform distribution of ferrite phase around the grains of the martensitic strengthening phase, and the grains of the martensite strengthening phase are mainly flaky.
  • the metallographic structure of the hot-dip galvanized Q&P steel is a three-phase structure of 45-75% martensite, 15-30% ferrite, and 10-25% austenite by volume fraction.
  • the hot-dip galvanized Q&P steel has a yield strength ⁇ 720MPa, a tensile strength ⁇ 1180MPa, an elongation rate ⁇ 19%, and a strong-plastic product ⁇ 23.0GPa%.
  • the hot-dip galvanized Q&P steel has a yield strength of 721-956 MPa, a tensile strength of 1184-1352 MPa, an elongation of 19-22.5%, and a strong-plastic product of 23.6-28.9 GPa%.
  • the austenite in the metallographic structure of the hot-dip galvanized Q&P steel has good thermal stability, the austenite transformation rate at -50°C is lower than 8%, and the austenite transformation rate at -190°C is lower than 30% .
  • the content of C is selected from the range of 0.17-0.23%, 0.19-0.21% and 0.18-0.21%.
  • the content of Si is selected from the range of 1.1-1.7%, 1.3-1.5%, 1.4-2.0% and 1.6-1.8%.
  • the content of Mn is selected from the range of 1.6-2.2%, 1.8-2.0%, 2.4-3.0% and 2.6-2.8%.
  • the content of Cr is ⁇ 0.35%, such as ⁇ 0.25%; the content of Mo is ⁇ 0.25%; the content of Nb is ⁇ 0.06%, such as ⁇ 0.04%; The content of Ti is ⁇ 0.065%, such as ⁇ 0.04%, such as 0.006-0.016%; the content of V is ⁇ 0.055%, such as ⁇ 0.035%.
  • the low-carbon low-alloy hot-dip galvanized Q&P steel with tensile strength ⁇ 1180 MPa is prepared by the following method:
  • the final rolling temperature of hot rolling is greater than or equal to Ar3, and then cooled to 550-680 °C for coiling;
  • the cold rolling reduction rate is 40-80%;
  • the cold-rolled steel plate is rapidly heated to 770-845 °C, and the rapid heating adopts one-stage or two-stage rapid heating; when one-stage rapid heating is used, the heating rate is 50-500 °C/s; when two-stage rapid heating is used, the 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 from 550-625°C at a heating rate of 30-500°C/s (eg 50-500°C/s). to 770 ⁇ 845°C; then soaking, soaking temperature: 770 ⁇ 845°C, soaking time: 10 ⁇ 60s;
  • hot-dip galvanizing After hot-dip galvanizing, rapidly cool to room temperature at a cooling rate of 30-150°C/s to obtain a hot-dip pure zinc GI product; or, after hot-dip galvanizing, heat to 480-550°C at a heating rate of 10-300°C/s Alloying treatment is carried out at °C, and the alloying treatment time is 5-20s; after the alloying treatment, it is rapidly cooled to room temperature at a cooling rate of 30-250°C/s to obtain an alloyed hot-dip galvanized GA product.
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • step 4 the whole process of rapid heat treatment and hot-dip galvanizing takes 43 to 186 s.
  • 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 845 °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 845 °C.
  • the cooling rate in the rapid cooling stage of the strip or steel plate is 50-150°C/s.
  • the chemical composition mass percentage of the low-carbon low-alloy hot-dip galvanized Q&P steel with tensile strength ⁇ 1180MPa is: C: 0.17-0.23%, Si: 1.1-1.7%, Mn: 1.6-2.2% , P ⁇ 0.015%, S ⁇ 0.005%, 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 metallographic structure of the hot-dip galvanized Q&P steel is a three-phase structure with a volume fraction of 45-75% martensite, 15-30% ferrite and 10-25% austenite, and its matrix The distribution of the structure is uniform, with obvious lamellar tempered martensite, and the grain size is 1-3 ⁇ m. There is a uniformly distributed ferrite phase around the grains of the martensitic strengthening phase. shaped organizational structure.
  • the hot-dip galvanized Q&P steel has a yield strength of 721-805 MPa, a tensile strength of 1184-1297 MPa, an elongation of 19.1-22.4%, and a strong-plastic product of 23.6-28 GPa%.
  • the austenite in the metallographic structure of the hot-dip galvanized Q&P steel has good thermal stability, the austenite transformation rate at -50°C is lower than 8%, and the austenite transformation rate at -190°C is lower than 30%.
  • the chemical composition mass percentage of the low-carbon low-alloy hot-dip galvanized Q&P steel with tensile strength ⁇ 1180MPa is: C: 0.16-0.23%, Si: 1.4-2.0%, Mn: 2.4-3.0% , Ti 0.006 ⁇ 0.016%, P ⁇ 0.015%, S ⁇ 0.002%, Al: 0.02 ⁇ 0.05%, can also contain one or two of Cr, Mo, Nb, V, and Cr+Mo+Ti+Nb +V ⁇ 0.5%, the balance is Fe and other inevitable impurities.
  • the low-carbon low-alloy hot-dip galvanized Q&P steel with a tensile strength of ⁇ 1180 MPa is a low-carbon low-alloy hot-dip galvanized Q&P steel with a tensile strength of ⁇ 1280 MPa.
  • the C content is 0.18-0.21%.
  • the Si content is 1.6-1.8%.
  • the Mn content is 2.6-2.8%.
  • the metallographic structure of the hot-dip galvanized Q&P steel is a three-phase structure of martensite, ferrite and austenite (the proportion of martensite structure is 75-90%, and the proportion of retained austenite structure is 10-25% %, the ferrite structure accounts for 3-10%), the matrix structure is evenly distributed, and there is obvious lamellar tempered martensite, the grain size is 1-3 ⁇ m, and there is a uniform martensite strengthening phase around the grains.
  • the distributed ferrite phase and martensitic strengthening phase grains are dominated by flaky structure.
  • the hot-dip galvanized Q&P steel has a yield strength of 802-956 MPa, a tensile strength of 1280-1352 MPa, an elongation of 19-22.5%, and a strong-plastic product of 25.2-28.9 GPa%.
  • the austenite in the metallographic structure of the hot-dip galvanized Q&P steel has good thermal stability, the austenite transformation rate at -50°C is lower than 8%, and the austenite transformation rate at -190°C is lower than 30%.
  • 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 exists in two phases in steel: ferrite and cementite. 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 martensite and pearlite will increase, which will greatly improve the strength and hardness of steel, but its plasticity and toughness will be obvious. 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 strengthening effect of the steel will also be significantly weakened, making the process performance of the steel worse, so the carbon content should be reduced as much as possible under the premise of ensuring the strength.
  • carbon is one of the most effective strengthening elements for the martensite matrix. It dissolves in austenite, expands the austenite phase, greatly improves the stability of austenite, and makes pearlite The transformation C curve of bainite and pearlite is shifted to the right, the transformation of pearlite and bainite is delayed, and the Ms point temperature is lowered. Too low carbon content will reduce the stability of retained austenite, and too high carbon content will cause twinning in martensite, reducing the plasticity, toughness and weldability of the steel. Considering comprehensively, the carbon content is limited within the range of 0.16-0.23%. In some embodiments, the content of C is 0.18-0.21%. In other embodiments, the content of C is 0.19-0.21%.
  • Mn Manganese can form a solid solution with iron, thereby improving the strength and hardness of ferrite and austenite in carbon steel, and enabling the steel to obtain finer and higher strength pearlite in the cooling process after hot rolling, and pearlite.
  • the content of Mn also increases with the increase of Mn content.
  • Manganese is also a carbide forming element, and manganese carbides can dissolve into cementite, thereby indirectly enhancing the strength of strengthening phases such as martensite and pearlite. Manganese can also strongly enhance the hardenability of steel, further increasing its strength.
  • the content of Mn is 1.8-2.0%. In other embodiments, the content of C is 2.6-2.8%.
  • the addition of manganese can reduce the martensitic transformation temperature Ms, increase the content of retained austenite, and improve the stability of retained austenite, and manganese has little effect on the toughness of the steel.
  • Ms martensitic transformation temperature
  • the manganese content is high, the grains of the steel tend to be coarsened, and the overheating sensitivity of the steel is increased.
  • the present invention limits the manganese content within the range of 1.6-3.0%.
  • Si Silicon forms a solid solution in ferrite or austenite, thereby enhancing the yield strength and tensile strength of steel, and silicon can increase the cold working deformation hardening rate of steel, which is a beneficial element in alloy steel.
  • silicon is obviously enriched on the intergranular fracture surface of silicon-manganese steel. The segregation of silicon at the grain boundary can slow down the distribution of carbon and phosphorus along the grain boundary, thereby improving the embrittlement state of the grain boundary. Silicon can improve the strength, hardness and wear resistance of steel, and it will not significantly reduce the plasticity of steel within a certain range. Silicon has strong deoxidation ability and is a commonly used deoxidizer in steelmaking. Silicon can also increase the fluidity of molten steel, so general steel contains silicon, but when the content of silicon in steel is too high, its plasticity and toughness will be reduced. Decreased significantly. For Q&P Steel:
  • silicon is a non-carbide forming element, and its solubility in carbides is extremely low. During the isothermal process of QP steel, it can inhibit the formation of Fe3C and enrich the untransformed austenite with carbon, thereby greatly improving the austenite. Stability, enabling it to be retained at room temperature;
  • silicon is a ferrite forming element, which can improve the stability of retained austenite and play a role in solid solution strengthening, thereby improving the strength of steel;
  • silicon has the effect of narrowing the austenite phase region and increasing the activity of C element in ferrite.
  • Silicon has no obvious effect on the growth rate of austenite, but has a significant effect on the shape and distribution of austenite.
  • the increase in silicon content will increase the difficulty of manufacturing in the pre-heat treatment process; the present invention limits the silicon content to 1.1-2.0% within the range.
  • the Si content is 1.3-1.5%. In other embodiments, the Si content is 1.6-1.8%.
  • 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 the steel.
  • chromium element When other alloying elements are added, the effect of chromium element to improve the strength and hardness of steel is more significant. Since Cr improves the quenching ability of steel during air cooling, it has an adverse effect on the weldability of steel. However, when the chromium content is less than 0.3%, the adverse effects on weldability can be ignored; when the content is greater than this, defects such as cracks and slag inclusions are likely to occur during welding. When Cr coexists with other alloying elements (such as coexisting with V), the adverse effect of Cr on weldability is greatly reduced. For example, when Cr, Mo, V and other elements exist in the steel at the same time, even if the Cr content reaches 1.7%, there is no significant adverse effect on the welding performance of the steel. In the present invention, chromium element is a beneficial and unnecessary addition element, and the addition amount should not be too much considering factors such as cost increase. In some embodiments, the Cr content is ⁇ 0.35%, such as ⁇ 0.25%.
  • Molybdenum suppresses the self-diffusion of iron and the diffusion rate of other elements.
  • the atomic radius of Mo is larger than that of ⁇ -Fe atoms.
  • Mo can increase the bond attraction of lattice atoms and increase the recrystallization temperature of ⁇ ferrite.
  • the strengthening effect of Mo in pearlitic, ferritic, martensitic steels, and even in high-alloy austenitic steels is also very obvious.
  • the good effect of Mo in steel also depends on the interaction with other alloying elements in the steel.
  • the solid solution strengthening effect of Mo is more significant. This is because when the strong carbide forming element is combined with C to form a stable carbide, it can promote the more effective dissolution of Mo into the solid solution, which is more conducive to the improvement of the thermal strength of the steel. Adding Mo can also increase the hardenability of steel, but the effect is not as significant as that of C and Cr. Mo will inhibit the transformation of the pearlite region and accelerate the transformation of the medium temperature region. Therefore, the steel containing Mo can form a certain amount of bainite and eliminate the formation of ferrite even when the cooling rate is high. One of the reasons why the thermal strength of alloy heat-resistant steel has a favorable effect.
  • Mo can also significantly reduce the hot brittle tendency of steel and reduce the speed of pearlite spheroidization.
  • Mo content is below 0.15%, there is no adverse effect on the weldability of the steel.
  • molybdenum element is a beneficial and unnecessary addition element, and the addition amount should not be too much considering factors such as cost increase.
  • the content of Mo is ⁇ 0.25%
  • Nb element is a carbide and nitride forming element, and can meet this requirement at relatively low concentration. 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. Adding Nb can prevent the growth of austenite grains and increase the coarsening temperature of steel grains. Nb element and carbon generate very stable NbC. Adding a small amount of Nb element to steel can use its precipitation strengthening effect to improve the strength of the matrix. Nb element has obvious hindering effect on the growth of ferrite recrystallization and austenite grain growth, which can refine the grain and improve the strength and toughness of steel; Transformation behavior and carbide formation also play a role.
  • Nb can increase the content of carbon in retained austenite, hinder the formation of bainite, promote the nucleation of martensite, obtain a dispersed martensite structure, and improve the stability of retained austenite.
  • Nb element to improve the strength of dual-phase steel can be used to obtain a certain strength of dual-phase steel under the condition of lower content of martensite and low C content, which can improve the strength and toughness of dual-phase steel; at the same time, adding Nb element has another benefit It is possible to increase the strength of steel in a wide annealing temperature range.
  • Nb 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 Nb is ⁇ 0.06%, such as ⁇ 0.04%.
  • Ti is a micro-alloying element and belongs to the ferrite forming element in the closed ⁇ region. It can increase the critical point of the steel. Ti and C in the steel can form a very stable TiC, which is generally within the austenitizing temperature range of heat treatment. Inside, TiC is extremely difficult to dissolve. Due to the grain refinement of austenite by TiC particles, when austenite decomposes and transforms, the opportunity for the formation of new phase nuclei increases, which accelerates the austenite transformation. In addition, Ti can form TiC and TiN precipitation phases with C and N, which are more stable than carbonitrides of Nb and V, significantly reduce the diffusion rate of C in austenite, and greatly reduce the formation rate of austenite.
  • the precipitated TiC has a precipitation strengthening effect; during the tempering process, Ti slows down the diffusion of C in the ⁇ phase, slows down the precipitation and growth of carbides such as Fe and Mn, increases the tempering stability, and can Plays a secondary hardening effect by precipitating TiC.
  • the high temperature strength of steel can be improved by microalloying of Ti.
  • Adding a small amount of Ti to steel can, on the one hand, increase the strength and improve the welding performance of the steel while reducing the carbon equivalent content; on the other hand, it can fix impure substances such as oxygen, nitrogen and sulfur, thereby improving the steel Weldability; secondly, due to the effect of its microscopic particles, such as the insolubility of TiN at high temperature, it can prevent the coarsening of grains in the heat-affected zone and improve the toughness of the heat-affected zone, thereby improving the weldability of steel.
  • the Ti content is ⁇ 0.065%, such as ⁇ 0.04%.
  • its addition amount may be in the range of 0.006-0.016%.
  • Microalloying element V is a ferrite stabilizing element and a strong carbide forming element, which has a strong grain refining effect and can make the steel structure dense.
  • the addition of V to the steel can improve the strength, plasticity and toughness of the steel at the same time.
  • Vanadium can also improve the high temperature strength of structural steel. Vanadium does not improve hardenability. Adding a small amount of microalloying element V to the steel can ensure that the steel can refine the grains through the dispersion and precipitation of carbon and nitride points (size less than 5nm) and the solid solution of V when the carbon equivalent of the steel is low.
  • the steel has good performance such as weldability.
  • Adding a small amount of V to the steel can, on the one hand, increase the strength and improve the welding performance of the steel while reducing the carbon equivalent content; Weldability; secondly, due to the effect of its microscopic particles, such as the insolubility of V(CN) at high temperature, it can prevent the coarsening of grains in the heat-affected zone and improve the toughness of the heat-affected zone, thereby improving the weldability of steel .
  • 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 V is ⁇ 0.055%, such as ⁇ 0.035%.
  • Nb, V, and Ti are the forming elements of carbides and nitrides. These elements can meet this requirement at relatively low concentrations.
  • Nb, V, Ti are strong carbide forming elements. It exists in the form of carbides, nitrides, and carbonitrides, and a small part is dissolved in ferrite.
  • the addition of these microalloying elements can strengthen the ferritic matrix through grain refinement and precipitation.
  • the formation of ferrite leads to carbon enrichment of retained austenite, which delays the transformation of austenite to bainite, while finely dispersed carbonitrides inhibit bainite nucleation, thereby also delaying the kinetics of 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 present 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 by means of a rapid heat treatment method (including rapid heating, short-time heat preservation and rapid cooling process), and finally obtains fine, uniform, dispersed Distribution of various organizational structures and good strong plastic matching.
  • a rapid heat treatment method including rapid heating, short-time 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.
  • Relatively high heating rates must be used to shorten the residence time of the structure in the high temperature range to ensure grain refinement.
  • the main phase structure of the Q&P steel obtained by the rapid heat treatment method of the present invention is martensite (volume fraction of 75-90%) and retained austenite (volume fraction of 10-25%), and contains a very small amount of Therefore, strictly speaking, its phase structure is multiphase structure, its matrix structure is evenly distributed, and there is obvious lamellar tempered martensite, and the grain size is 1 ⁇ 3 ⁇ m, the grains of martensite strengthening phase are surrounded by uniformly distributed ferrite phase, and the grains of martensite strengthening phase are mainly flaky structure.
  • the manufacturing method of the low-carbon low-alloy Q&P steel with tensile strength ⁇ 1180MPa comprises the following steps:
  • the cold rolling reduction rate is 40% to 85%, and the rolled hard strip or steel plate is obtained;
  • the strip or steel plate after cold rolling is rapidly heated to 770 ⁇ 845°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 During the heating process, 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 from 550-625°C to 770-845°C at a heating rate of 50-500°C/s;
  • Soaking is carried out at the target temperature of 770-845 °C in the two-phase region of austenite and ferrite, and the soaking time is 10-60s;
  • the strip or steel plate After soaking, the strip or steel plate is slowly cooled to 700-770°C (eg 720-770°C) at a cooling rate of 5-15°C/s, and then cooled to 50-200°C/s (eg 50-150°C/s) The cooling rate is rapidly cooled to 230-280°C, and the temperature is kept at this temperature range for 2-10s;
  • 700-770°C eg 720-770°C
  • 50-200°C/s eg 50-150°C/s
  • the strip or steel plate is heated to 300-470°C at a heating rate of 10-30°C/s for tempering, and the tempering time is 10-60s
  • the strip or steel plate After tempering, the strip or steel plate is cooled to room temperature, and the cooling rate is 30-100°C/s.
  • the final rolling temperature of the hot rolling is ⁇ Ar3.
  • the coiling temperature is 580-650°C.
  • the cold rolling reduction ratio is 60-80%.
  • the entire process of the rapid heat treatment takes 71 to 186 s.
  • 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; 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 845 °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; 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 845 °C.
  • the final temperature of the rapid heating is 790-845°C.
  • the rapid cooling rate of the strip or steel plate is 50-150°C/s.
  • the steel strip or steel plate is slightly heated or cooled within the soaking time period, the temperature after heating is not more than 845°C, and the temperature after cooling is not lower than 770°C.
  • the soaking time is 10-40s.
  • the rapid heat treatment and hot-dip galvanizing manufacturing method of low-carbon and low-alloy hot-dip galvanized Q&P steel with tensile strength ⁇ 1180 MPa comprises the following steps:
  • the final rolling temperature of hot rolling is greater than or equal to Ar3, and then 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 ⁇ 845°C;
  • Soaking is carried out at the target temperature of 770-845 °C in the two-phase region of austenite and ferrite, and the soaking time is 10-60s;
  • the strip or steel plate After soaking, the strip or steel plate is cooled slowly to 720-770°C at a cooling rate of 5-15°C/s; then rapidly cooled to 230-280°C at a cooling rate of 50-200°C/s (eg 50-150°C/s). °C, and keep it in this temperature range for 2 to 10s (such as 2 to 8s);
  • the strip or steel plate is heated to 460-470°C at a heating rate of 10-30°C/s for distribution treatment, and the distribution time is 10-60s;
  • the strip or steel plate is dipped into the zinc pot for hot-dip galvanizing;
  • the whole process of the rapid heat treatment and hot-dip galvanizing takes 43-186 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 845 °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 845 °C.
  • the final temperature of the rapid heating is 790-845°C.
  • the cooling rate in the rapid cooling stage is 50-150°C/s.
  • the strip or steel plate is heated or cooled in a small range during the soaking time period, and the temperature after heating is not more than 845°C, and the temperature after cooling is not lower than 770°C.
  • the soaking time is 10-40s.
  • 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 can be quantitatively described by the relationship affected by the heating rate.
  • the functional relationship between the ferrite recrystallization volume fraction and the temperature T during the continuous heating process is:
  • X(t) is the volume fraction of ferrite recrystallization
  • n is the Avrami index, which is related to the phase transformation mechanism and depends on the decay cycle of the recrystallization nucleation rate, generally in the range of 1 to 4
  • T is Heat treatment temperature
  • T star is the recrystallization start temperature
  • ⁇ is the heating rate
  • b(T) is obtained by the following formula:
  • the deformed matrix first recovers, recrystallizes and grains grow, and then the phase transformation from ferrite to austenite occurs, and the phase deformation nuclei are mainly concentrated in the grown ferrite crystals. At the boundary, the nucleation rate is low, so the final grain structure is relatively coarse.
  • the deformed matrix begins to recrystallize before it can fully recover, and the phase transformation of ferrite to austenite begins before the recrystallization is completed or the grain growth begins.
  • the grains are small and the grain boundary area is large, so the nucleation rate is significantly improved, and the grains are obviously refined.
  • a large number of nucleation points are provided for austenite due to the retention of a large number of crystal defects such as dislocations in the ferrite crystal, which makes austenite appear Explosive nucleation, further refinement of austenite grains.
  • the retained high-density dislocation line defects also become the channels for the high-speed diffusion of carbon atoms, so that each austenite grain can rapidly grow and grow, thus increasing the austenite volume fraction.
  • the present invention sets the heating rate as 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 impact on the recovery process, and the heating rate is controlled to be 15-300°C/s, more preferably 50-300°C/s; the heating temperature is from 500-625°C When the austenitizing temperature reaches 770-845°C, the heating rate has the greatest influence on the recrystallization nucleation, phase deformation nucleation and grain growth process, and the heating rate is controlled to be 50-300°C/s; 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 depends on the C content. In the traditional process, the soaking temperature is generally set at 30 to 50°C above A C3 .
  • the present invention uses the rapid heating technology to form a large number of dislocations in the ferrite, which provides a shape for the austenite transformation. Nuclear work, so just heat the temperature to between A C1 and A C3 .
  • the C content in the Q&P steel of the present invention is 0.16-0.23%, and A C1 and A C3 are respectively about 730°C and 870°C.
  • Q&P steel contains a large number of undissolved and uniformly distributed carbides, which can mechanically hinder the growth of austenite grains during the soaking process, which is beneficial to refine the grain size of alloy steels, but If the soaking temperature is too high, the number of undissolved carbides will be greatly reduced, which will weaken this hindering effect, enhance the growth tendency of grains, and then reduce the strength of the steel. When the amount of undissolved carbide is too large, it may cause aggregation, resulting in uneven distribution of local chemical components. When the carbon content at the aggregation is too high, local overheating will also occur.
  • 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 a higher volume fraction and a uniform and fine martensite structure after cooling. Excessive soaking temperature will make the austenite grains coarse, the workpiece is easy to crack during the quenching process, and the martensite structure obtained after quenching will also be coarser, resulting in poor mechanical properties of the steel; it will also reduce the retained austenite. , reduce the hardness and wear resistance of the workpiece. If the soaking temperature is too low, the content of carbon and alloying elements dissolved in the austenite will be insufficient, so that the concentration of alloying elements in the austenite will be unevenly distributed, and the hardenability of the steel will be greatly reduced. performance is adversely affected.
  • the soaking temperature of hypoeutectoid steel should be Ac 3 +30 ⁇ 50°C.
  • the existence of carbide-forming elements will hinder the transformation of carbides, so the soaking temperature can be appropriately increased.
  • the present invention selects 770-845° C. as the soaking temperature, in order to obtain a reasonable quenching process and ideal microstructure and properties.
  • 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;
  • the shorter the heating 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 soaking time is too short, the distribution of alloying elements in the steel will be uneven. uniform, resulting in insufficient austenitization; too long soaking time will easily lead to coarse austenite grains.
  • the length of soaking time is also related to the content of carbon and 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 10-60s.
  • the cooling rate of the sample during rapid cooling must be greater than the critical cooling rate to obtain the martensite structure.
  • the critical cooling rate mainly depends on the material composition.
  • the Si content is 1.1-2.0%, and the Mn content It is 1.6-3.0%, and the content is relatively high, so Si and Mn greatly enhance the hardenability of Q&P 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.
  • the rapid cooling rate is set to 50-200°C/s.
  • the alloy steel when the alloy steel is tempered below 150°C, the alloy elements cannot diffuse due to the low temperature, and only the carbon element has a certain diffusion ability. Therefore, although the low temperature tempered steel has high hardness, its brittleness is too large. , the toughness is very poor and cannot meet the performance requirements of the workpiece.
  • tempering is performed at a temperature above 200 °C, carbon and other alloying elements contained in martensite will begin to precipitate in large quantities, reducing residual stress until it disappears, and the hardness of tempered steel will gradually increase with the increase of tempering temperature. decreased, but increased toughness.
  • the tempering temperature reaches about 500 °C
  • the decomposition of martensite ends, the cementite gradually aggregates and grows, the ⁇ phase begins to undergo a recovery process, and the temperature continues to increase, and the ⁇ phase begins to recrystallize to form polygonal ferrite. Decreased significantly.
  • the ultimate purpose of the present invention is to obtain better strength and plasticity at the same time, so the present invention sets the tempering temperature at 300 ⁇ 470°C.
  • the tempering time plays three roles: (1) to ensure the sufficient transformation of the structure; (2) to reduce or eliminate the internal stress; (3) to cooperate with the tempering temperature to obtain the required properties of the workpiece .
  • the rapid heating technology is used to refine the austenite grains, thereby shortening the distance between the retained austenite and martensite generated after a rapid cooling, and the carbon atoms change from supersaturated martensite to retained austenite.
  • the efficiency of bulk diffusion partitioning is increased, so the time required for the tempering process is also greatly reduced.
  • the present invention sets the tempering time to 10-60s.
  • the Q&P steel with higher alloy content is carbon partitioned (tempered) below 150 °C. Because the temperature is too low, the alloying elements cannot be diffused, and only the carbon element has a certain diffusion ability. Therefore, although the low temperature partition steel has higher Hardness, but its brittleness is too large and its toughness is poor, which cannot meet the performance requirements of the workpiece.
  • partitioning is performed at a temperature above 200 °C, carbon and other alloying elements contained in martensite will begin to precipitate in large quantities, reducing the residual stress until it disappears, and the hardness of the partitioned steel will gradually decrease with the increase of the partitioning temperature.
  • the partition temperature reaches about 500°C
  • the decomposition of martensite ends, the cementite gradually aggregates and grows, and the ⁇ phase begins to undergo a recovery process.
  • the ⁇ phase begins to recrystallize to form polygonal ferrite.
  • the main purpose of the partitioning process of the present invention is to diffuse the carbon in the martensite that has been obtained into the martensite that has not yet occurred.
  • the carbon in the martensite is reduced and the plasticity is improved, while the carbon concentration diffused into the retained austenite is increased, and its stability is enhanced, so that the final product can obtain better strength and plasticity at the same time.
  • the distribution temperature is set at 460 ⁇ 470 °C.
  • the distribution time plays three roles: (1) to ensure that the structural transformation is fully carried out; (2) to reduce or eliminate the internal stress as much as possible; (3) to cooperate with the distribution temperature to obtain the required properties of the workpiece.
  • the rapid heating technology is used to refine the austenite grains, thereby shortening the distance between the retained austenite and martensite generated after one rapid cooling, and the carbon atoms change from supersaturated martensite to retained austenite.
  • the efficiency of diffusive partitioning is increased, so the time required for the partitioning process is also greatly reduced.
  • the distribution time is too short, it is difficult to eliminate the internal stress and reduce the brittleness of the workpiece.
  • the invention sets the distribution time at 10-60s.
  • 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 alloy content in the same grade of steel can be reduced, the crystal grains can be refined, the good soft and hard phase structure and the matching of strength and toughness can be obtained;
  • the rapid cooling process is transformed to realize the rapid heat treatment process, which can greatly shorten the length of the heating and soaking section of the annealing furnace (at least one third shorter than the traditional continuous annealing furnace), and improve the production of the traditional continuous hot-dip galvanizing unit. efficiency, reduce production costs and energy consumption, significantly reduce the number of continuous annealing furnace rolls, especially the number of high-temperature furnace rolls, which can reduce energy consumption and equipment investment.
  • 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 metallographic structure of the obtained Q&P steel is 75-90% martensite, 10-25% retained austenite, and 3% ferrite.
  • the metallographic structure of the obtained hot-dip galvanized Q&P steel is a refined three-phase structure of martensite, ferrite and austenite, and the volume fraction of martensite is preferably 45-75% %, the volume fraction of retained austenite is 10-25%, and the volume fraction of ferrite is 15-30%.
  • the matrix structure of the obtained Q&P steel and hot-dip galvanized Q&P steel is uniformly distributed, with obvious lamellar tempered martensite, and the grain size is refined to 1-3 ⁇ m, and there is a uniform distribution around the martensite-strengthening phase grains.
  • the ferrite phase is mainly in the form of ferrite, and the grains of the martensite strengthening phase are mainly flaky structure; the austenite in the structure has various forms such as block, strip and granular, and has good thermal stability, -50
  • the °C austenite transformation rate is lower than 8%, and the -190 °C austenite transformation rate is lower than 30%, and the TRIP effect can continue to occur under different strain conditions, so the product mechanical properties and user performance are excellent.
  • the alloy composition of the Q&P steel obtained by the present invention is greatly reduced, the grain size is reduced by 40-80%, and the performance is excellent; the yield strength is 668-1112MPa, and the tensile strength is 1181 ⁇ 1350MPa, the elongation is 18.9 ⁇ 24.2%, and the strong-plastic product is 24.1 ⁇ 28.97GPa%.
  • the average grain size of the Q&P steel obtained by the rapid heat treatment of the present invention is 1-3 ⁇ m under the premise that the manufacturing conditions of the previous process remain unchanged.
  • the grain size is reduced by 10-40%, and a good effect of fine-grain strengthening can be obtained; the yield strength is greater than or equal to 720MPa, the tensile strength is greater than or equal to 1180MPa, the elongation is greater than or equal to 19%, and the strong-plastic product is greater than or equal to 23.0GPa%; preferably, all
  • the yield strength of the hot-dip galvanized Q&P steel is 721-956MPa, the tensile strength is 1184-1352MPa, the elongation is 19-22.5%, and the strength-plastic product is 23.6-28.9GPa%.
  • the rapid heat treatment process of low-carbon and low-alloy Q&P steel with tensile strength ⁇ 1180MPa and the rapid heat-treatment process of low-carbon and low-alloy hot-dip galvanized Q&P steel according to the present invention can be shortened to 71-186s and 43 seconds in the whole process of heat treatment. ⁇ 186s, greatly reducing the time of the entire heat treatment process (the traditional continuous annealing process time is usually 5-8min), significantly improving production efficiency, reducing energy consumption, and reducing production costs.
  • the rapid heat treatment method of the present invention shortens the heating section and soaking section time by 60-80%, and the entire heat treatment process time is shortened to 71-186s;
  • Hot-dip galvanized Q&P steel and its heat treatment process the rapid heat treatment method of the present invention shortens the length and time of the heating section and soaking section of the continuous hot-dip galvanizing annealing furnace (compared with the traditional continuous hot-dip galvanizing annealing furnace, the heating section and the soaking section are shorter than the traditional continuous hot-dip galvanizing annealing furnace.
  • the length of the hot section can be shortened by 60-80%) and the entire heat treatment process time.
  • the invention can save energy, reduce emissions, reduce consumption, significantly reduce the one-time investment in furnaces and other equipment, and 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, reduce heat treatment and The production cost of the previous process reduces the manufacturing difficulty of each process before the heat treatment.
  • the rapid heat treatment technology can reduce the heating process and soaking process time, shorten the length of the furnace, reduce the number of furnace rolls, and make the furnace The probability of generating surface defects is reduced, and the surface quality of the product will be significantly improved; in addition, due to the refinement of the product grain and the reduction of the material alloy content, the Q&P steel obtained by the technology of the present invention has processing and forming properties such as hole expanding properties and bending properties, welding User performance such as performance has also improved.
  • 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 low-carbon and low-alloy Q&P steel with tensile strength ⁇ 1180MPa obtained by the invention has important value for the development of new-generation light-weight vehicles, trains, ships, airplanes and other transportation vehicles and the healthy development of corresponding industries and advanced manufacturing industries.
  • Fig. 1 is the microstructure picture of Q&P steel produced according to Example 1 of Test Steel A in Example 1 of the present invention.
  • Fig. 2 is the microstructure picture of the Q&P steel produced by the traditional process 1 of the test steel A in the first embodiment of the present invention.
  • Example 3 is a picture of the microstructure of the Q&P steel produced in Example 7 of the test steel K in Example 1 of the present invention.
  • Example 4 is a picture of the microstructure of the Q&P steel produced in Example 8 of the test steel R according to Example 1 of the present invention.
  • FIG. 5 is a picture of the microstructure of the Q&P steel produced according to Example 22 of the test steel P in Example 1 of the present invention.
  • Fig. 6 is the microstructure picture of Q&P steel produced according to Example 23 of Test Steel S in Example 1 of the present invention.
  • Fig. 8 is the microstructure picture of the Q&P steel produced by the traditional process 1 of the test steel A in the second embodiment of the present invention.
  • Example 10 is a picture of the microstructure of the Q&P steel produced in Example 8 of the second test steel R of the present invention.
  • Example 11 is a picture of the microstructure of the Q&P steel produced in Example 22 of the second test steel P of the present invention.
  • Example 12 is a picture of the microstructure of the Q&P steel produced in Example 23 of the second test steel S of the present invention.
  • Example 13 is a microstructure picture of the hot-dip pure galvanized Q&P 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 dual-phase steel (GA) produced according to Example 17 of Test Steel I in Example 3 of the present invention.
  • Fig. 16 is the microstructure picture of the hot-dip pure zinc dual-phase 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 dual-phase 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 Q&P steel (GI) produced in Example 1 of the test steel A in Example 4 of the present invention.
  • FIG. 19 is a microstructure picture of the hot-dip pure zinc Q&P steel (GI) produced by the traditional process 1 of the fourth test steel A of the present invention.
  • Example 20 is a microstructure picture of the alloyed hot-dip galvanized dual-phase 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 dual-phase 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 dual-phase 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 Tensile Test of Metallic Materials Part 1: Test Method at Room Temperature", and the P7 sample was used to test in the transverse direction.
  • Table 1 for the composition of the test steel in this embodiment, refer to Table 2 and Table 3 for the specific parameters of the present embodiment and the traditional process, and Table 4 and Table 5 for the composition of the test steel of this embodiment. 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 Q&P steel obtained by the method of the present invention is 668-1002MPa, the tensile strength is 1181-1296MPa, the elongation is 18.9-24.2%, and the strong-plastic product is 24.1-28.6GPa%.
  • 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 material structure after different heat treatment methods.
  • the microstructure (Fig. 1) of steel A treated by the embodiment of the present invention is mainly composed of fine and uniform austenite structure and a small amount of carbides dispersed on the martensite matrix, austenite and martensite grains.
  • the structure and carbides are very fine and evenly distributed in the matrix, which is very beneficial to improve the strength and plasticity of the material.
  • the structure of steel A treated by the traditional process Fig.
  • FIG. 3 is a microstructure diagram of typical composition K steel obtained by Example 7, and FIG. 4 is a microstructure diagram of typical composition R steel obtained by Example 8.
  • FIG. 5 is a microstructure diagram of typical composition P steel obtained by Example 22, and
  • FIG. 6 is a microstructure diagram of typical composition S steel obtained by Example 23.
  • Embodiments 7, 8, 22, and 23 are all processes with a shorter entire heat treatment cycle. It can be seen from the figure that by using the method of the present invention, a more uniform, fine and dispersed phase structure can be obtained after a short-time rapid annealing treatment. Therefore, the preparation method of the present invention can refine the crystal grains, so that the structure of each phase of the material is evenly distributed in the matrix, thereby improving the structure of the material and improving the performance of the material.
  • composition of the test steel of the present invention is shown in Table 6, the specific parameters of the embodiment of the present invention and the traditional process are shown in Table 7 and Table 8, and Table 9 and Table 10 are the main properties of the steel prepared by the test steel composition of the present invention according to the embodiment and the traditional process. .
  • 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 between the material structure and strength and toughness.
  • the Q&P steel obtained by the method of the present invention has a yield strength of 754-1112 MPa, a tensile strength of 1281-1350 MPa, an elongation of 19-22.2%, and a strong-plastic product of 24.8-28.97 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 material structure after different heat treatment methods.
  • the microstructure of the obtained steel treated in the embodiment of the present invention is mainly composed of fine and uniform martensite structure and a small amount of carbides dispersed on the ferrite matrix.
  • the martensite grain structure and a small amount of carbides are very small and It is evenly distributed in the ferrite matrix, which is very beneficial to improve the strength and plasticity of the material.
  • the distribution of the steel structure treated by the traditional process is relatively uneven, the martensite grains are relatively large, and a small amount of retained austenite and carbide structures are distributed on the martensite grain boundaries, and the distribution is uneven.
  • the characteristics of the structure treated by the traditional process are: the grains are relatively coarse, and there is a certain uneven distribution of the structure.
  • FIG. 9 is a microstructure diagram of typical composition K steel obtained by Example 7, and FIG. 10 is a microstructure diagram of typical composition R steel obtained by Example 8.
  • FIG. 11 is a microstructure diagram of typical composition P steel obtained by Example 22, and
  • FIG. 12 is a microstructure diagram of typical composition S steel obtained by Example 23.
  • Embodiments 7, 8, 22, and 23 are all processes with a shorter entire heat treatment cycle. It can be seen from the figure that by using the method of the present invention, a more uniform, fine and dispersed phase structure can be obtained after a short-time rapid annealing treatment. Therefore, the preparation method of the present invention can refine the crystal grains, so that the structure of each phase of the material is evenly distributed in the matrix, thereby improving the structure of the material and improving the performance of the material.
  • the traditional continuous annealing unit can be transformed by adopting the rapid heating and rapid cooling process to realize the rapid heat treatment process, which can greatly shorten the length of the heating and soaking section of the traditional continuous annealing furnace , to improve the production efficiency of the traditional continuous annealing unit, 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 high surface quality strip products;
  • the new continuous annealing unit with rapid heat treatment 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 strip steel grains, further improve the material strength, and reduce the alloy cost. And the manufacturing cost and manufacturing difficulty of the pre-heat treatment process, and improve the user performance such as the welding performance of the material.
  • 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 is about 10 per minute, which can significantly improve the surface quality of the strip.
  • the rapid heat treatment process of recrystallization and austenitization completed in a very short time will also provide a more flexible and flexible high-strength steel structure design method, and then without changing the alloy composition and rolling process, etc.
  • the material structure can be improved and the material properties can be improved.
  • Advanced high-strength steel represented by Q&P steel has 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 Q&P steel.
  • Table 11 for the composition of the test steel in this example, and see Table 12 (one-stage heating) and Table 13 (two-stage heating) for the specific parameters of the embodiment of the present invention and the traditional process;
  • Table 14 and Table 15 are the composition of the test steel of the present invention Main properties of the obtained GI and GA hot-dip galvanized QP steel products prepared according to the examples and traditional 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 Q&P steel obtained by the method of the present invention is 721-805MPa
  • the tensile strength is 1184-1297MPa
  • the elongation is 19.1-22.4%
  • the strong-plastic product is 23.6-28GPa%.
  • FIG. 13 and 14 are microstructure diagrams of typical composition A steel through Example 1 and Conventional Process Example 1. From the figure, there is a very big difference in the structure after hot-dip galvanizing.
  • the microstructure of the steel A after the rapid heat treatment of the present invention (Fig. 13): the matrix microstructure is uniformly distributed, the microstructure has obvious lamellar tempered martensite, and the grain size is 1-3 ⁇ m. There is a uniformly distributed ferrite phase around the martensitic strengthening phase grains. Due to the decrease in the stability of martensite formed after part of the prior austenite grows, a small amount of tempered martensite appears in the microstructure after heat treatment, and the remaining martensite strengthening phase is still dominated by flaky morphology. The intenite grain structure and carbides are very fine and evenly distributed in the matrix, which is very beneficial to improve the strength and plasticity of the material.
  • the structure of the steel treated by the traditional process (Fig. 14) is a typical Q&P steel structure.
  • the lath martensite has coarse grains, austenite and carbides are distributed along the martensite grain boundaries, and the multiphase structure is unevenly distributed. .
  • 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).
  • Examples 17, 22, and 34 are all processes with a short heat treatment cycle; it can be seen from the figure that by using the method of the present invention, very uniform, fine, and dispersed phase structures can be obtained. Therefore, the preparation method of the hot-dip galvanized Q&P 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 The composition of the test steel in this example is shown in Table 16, and the specific parameters of the embodiment of the present invention and the traditional process are shown in Table 17 (one-stage heating) and Table 18 (two-stage heating); Table 19 and Table 20 are the composition of the test steel of the present invention.
  • Table 19 shows the test steel composition of the present invention according to the examples in Table 17 and Table 18 and the traditional process to prepare the obtained GI and GA hot-dip galvanized QP steel products. main performance.
  • the Q&P steel obtained by the method of the invention has a yield strength of 802-956 MPa, a tensile strength of 1280-1352 MPa, a maximum elongation of 19-22.5%, and a strong-plastic product of 25.2-28.9 GPa%.
  • FIG. 18 and 19 are microstructure diagrams of typical composition A steel through Example 1 and Conventional Process Example 1. From the figure, there is a very big difference in the structure after hot-dip galvanizing.
  • the microstructure of steel A after the rapid heat treatment of the present invention (Fig. 18) is composed of martensite, austenite and a small amount of ferrite and carbide. body, the grain size is 1-3 ⁇ m. Most of the strengthening phase grains are surrounded by ferrite. Due to the decrease in the stability of martensite formed after part of prior austenite grows, a small amount of tempered martensite appears in the microstructure after heat treatment, and the remaining strengthening phase is still dominated by massive morphology, ferrite and martensite crystals. The granular structure and carbides are very fine and evenly distributed in the matrix, which is very beneficial to improve the strength and plasticity of the material.
  • the structure of the steel treated by the traditional process (Fig. 19) is a typical Q&P steel structure.
  • the lath martensite has coarse grains, austenite and carbides are distributed along the martensite grain boundary, and the multiphase structure is unevenly distributed. .
  • 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).
  • Examples 17, 22, and 34 are all processes with a short heat treatment cycle; it can be seen from the figure that by using the method of the present invention, very uniform, fine, and dispersed phase structures can be obtained. Therefore, the preparation method of the hot-dip galvanized Q&P 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 invention transforms the traditional continuous annealing hot-dip galvanizing unit by adopting the rapid heating and rapid cooling process, so that the rapid heat treatment and hot-dip galvanizing process can be realized, and the heating section and soaking section of the traditional continuous annealing hot-dip galvanizing furnace can be greatly shortened. It can improve the production efficiency of traditional continuous annealing hot-dip galvanizing units, reduce production costs and energy consumption, reduce the number of furnace rolls in continuous annealing hot-dip galvanizing furnaces, and significantly reduce surface defects such as roll marks, pitting, and scratches.
  • the hot-dip galvanizing unit can be short and compact, flexible in material transition, It has the advantages of strong control ability and so on; for the material, it can refine the strip steel grain, further improve the material strength, reduce the alloy cost and the manufacturing difficulty of the pre-heat treatment process, and improve the user performance of the material such as forming and welding.
  • the present invention greatly promotes the technological progress of continuous annealing and hot-dip galvanizing of cold-rolled strip steel by adopting rapid heat treatment and hot-dip galvanizing process.
  • the intensification 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 and hot-dip galvanizing.
  • the number of rollers in the furnace of the unit should not exceed 10 rollers in the high temperature furnace section of the rapid heat treatment and hot-dip galvanizing production line with a speed of about 180 m/min, which can significantly improve the surface quality of the strip.
  • the rapid heat treatment and hot-dip galvanizing process of recrystallization and austenitization completed in a very short time will also provide a more flexible and flexible structure design method of high-strength steel, and then do not need to change the alloy composition and rolling. Under the premise of pre-process conditions such as manufacturing process, the material structure can be improved and the material properties can be improved.
  • the advanced high-strength steel represented by hot-dip galvanized Q&P steel has broad application prospects, and rapid heat treatment and hot-dip galvanizing technology have great development value.
  • the combination of the two will surely be the development and production of hot-dip galvanized Q&P steel. Provide more space.

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Abstract

抗拉强度≥1180MPa的低碳低合金Q&P钢或热镀锌Q&P钢及其制造方法。其化学成分质量百分比为：C：0.16～0.23％，Si：1.1～2.0％，Mn：1.6～3.0％，P≤0.015％，S≤0.005％，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。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工序。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，通过抑制铁素体再结晶得到等轴细晶复相的微观组织，最终获得的钢的金相组织为马氏体、残余奥氏体、铁素体的多相组织，晶粒尺寸在1～3μm，力学性能得以优化提高，材料性能区间范围得以扩展。

Description

抗拉强度≥1180MPa的低碳低合金Q＆P钢或热镀锌Q＆P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≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P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，这也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，以Q&P(Quenching and Partitioning，淬火及碳的再分配)钢为代表的第三代先进高强钢的开发日益得到重视。
Q&P热处理工艺是由Speer等人于21世纪初提出的一种新型的连续热处理工艺技术，该工艺主要包括四步：
第一，将带钢加热到奥氏体化温度保温；
第二，将试样快速冷却到M s～M f之间的某一温度，得到主要为马氏体和残余奥氏体的双相组织；
第三、将带钢加热至不高于M s的温度下保温，使碳元素从过饱和的马氏体向奥氏体中扩散配分，降低马氏体中的碳含量和硬度，改善其塑性，同时提高奥氏体的碳含量并增加其稳定性；
第四、冷却到室温，在这一个过程中，如果残余奥氏体的稳定性不够，部分奥氏体将转变为马氏体，那么室温下得到的残余奥氏体量将减少。
Q&P钢本质上是一种马氏体钢，但是它区别于传统的回火马氏体钢，在与回火马氏体钢同等强度下，Q&P钢的塑性却有很大的提高。这是由于Q&P钢的组织中存在残余奥氏体，这部分奥氏体在变形过程中转变为马氏体，产生所谓的TRIP效应，大幅度提高了钢的塑性。
目前，针对Q&P工艺的开发手段有两条，一是通过添加合金元素，提高钢中合金元素对碳化物析出的抑制能力；二是优化工艺，找出最佳温度和时间，通过调整Q&P工艺中淬火及配分过程的温度及时间，来改变Q&P钢的组织性能。
美国专利申请US2003/027825提出了Q&P钢生产工艺的大致过程，并将奥氏体化过 程限定于在高温下进行，材料组织需全奥氏体化，对于实际生产过程这一温度过高(850-950℃)，且时间长(通常要求钢板奥氏体化过程需保温2～5min)，设备要求较高，制造成本也较高。
中国专利CN1081931138B公开了“一种980MPa级汽车用冷轧高强Q&P钢及其生产方法”，该钢化学成分质量百分比为：C：0.18～0.24％，Si：0.6～1.3％，Mn：1.6～2.4％，P：0.02～0.04％，S≤0.005％，Nb：0.04～0.07％，N≤0.006％，Als：0.5～1.0％，余量为Fe和其它不可避免的杂质元素。其热轧工序终轧温度870～910℃，卷取温度660～710℃；所述冷轧工序的冷轧压下率≥45％；所述连续退火工序的均热段保温温度770～840℃，过时效段保温温度300～440℃，均热段保温时间为60～225s，过时效段保温时间为300～1225s；所述平整工序平整延伸率为0.3～0.9％。所得到的钢板的屈服强度大于550MPa，抗拉强度大于980MPa，断后伸长率大于18％。
该发明钢的主要特征在于通过传统Q&P工艺，获得良好强塑性配合的结果。由于采用传统热处理工艺方法，其均热时间和配分时间均很长，同时其合金含量相对较高，这也会增加制造成本、降低制造灵活性。
中国专利申请CN109136779A公开了“一种马氏体基体1100MPa级稀土Q&P钢制备方法”，该发明钢化学成分质量百分比为：C：0.15～0.22％，Si：0.6～1.7％，Mn：1.1～2.4％，Mo：0.1～0.5％，Al：0.1～0.5％，V：0.05～0.11％，Y：0.01～0.05％，P：0.02～0.04％，S≤0.005％，Nb：0.04～0.07％，N≤0.006％，B：0.001～0.006％，余量为Fe和其它不可避免的杂质元素。所得钢板的抗拉强度在1100MPa左右，断后伸长率20％左右。
该发明钢的主要特征在于通过同时添加稀土Y和Mo、V、Nb等合金元素用于细化晶粒，减少Mn元素的含量提升焊接性能，其制造过程需要进行两次铸造。冶炼过程：按照该发明所给出的成分配方，配料后经过转炉冶炼，真空炉二次精炼、铸造得到铸坯；微量元素熔入过程：在电弧熔炼炉中加入微量合金元素粉末(Mo、Al、V、Y、Nb、N，B等)，得到二次铸坯。热轧过程：使用加热炉将铸坯加热到1100-1150℃并进行保温1-3h，随后进行热轧，终轧温度为820-880℃，卷取温度550-650℃，所获得的钢板厚度为1.5-3.0mm，后水淬到室温；冷轧过程：酸洗后进行多道次冷轧，获得厚度为1.2-1.5mm的钢板；
其整个退火过程分为三次：
第一次：两相区锰配分过程：将材料以10-30℃/s加热到A C3和A C1(双相区)之间某一温度并保温3-15min后水淬至室温；
第二次：一次碳配分过程：将材料在M S与M f之间某一温度T 0进行保温10-300s，后将材料水淬至室温；
第三次：二次碳配分过程：将材料在M S与M f之间某一温度T 1(T 1温度比T 0稍低)进行10-300s的二次碳分配，后将材料水淬至室温。
该发明制造工艺复杂，能耗高，合金含量高且复杂，且多次水淬处理，涉及到材料表面氧化层的去除，带来环境和能耗等方面的诸多问题，导致制造成本增加和制造灵活性的降低。
中国专利申请CN108431248A公开了“一种用于制造具有改善延展性和可成形性的高强度钢板的方法和所获得的钢板”，该发明钢化学成分质量百分比为：C：0.15～0.23％，Mn：2.0～2.8％，Si：1.0～2.1％，Al：0.02～1.0％，Al+Si：1.0～2.1％，Nb：0～0.035％，Mo≤0.3％，Cr≤0.04％，余量为Fe和其它不可避免的杂质元素。所述钢板在退火温度TA下退火以获得包含至少65％的奥氏体和最多至35％的铁素体的组织。使所述钢板以至少20℃/秒的冷却速率从至少600℃的温度淬火至Ms-170℃到Ms-80℃的淬火温度QT，将所述钢板加热至350℃至450℃的配分温度PT，并将所述钢板保持在该配分温度下80-440s的配分时间Pt，随后立刻将所述钢板冷却至室温。所得钢板的抗拉强度大于1180MPa，断后伸长率大于12％。
该发明钢的主要特征在于采用高Mn、高Si和高Al成分通过传统Q&P工艺，控制最终组织中各相的比例，达到获得良好强塑性配合的结果。由于采用传统热处理工艺方法，其均热时间和配分时间均很长，这会增加制造成本、降低制造灵活性。
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。该发明钢的主要制造步骤如下：
1)将一定化学成分的铸坯锻造成锻坯，重新加热后热轧，水冷后卷曲，得到热轧带钢
2)将所述热轧带钢经酸洗后冷轧成冷轧带钢；
3)将所述冷轧带钢完全奥氏体化后保温一段时间，然后水冷到室温，生成全马氏体组织的预淬火带钢；
4)将所述预淬火带钢表面除鳞，从而去除氧化铁皮层和脱碳层，之后重新加热退火保温一段时间，然后用盐浴冷却到一定温度，经过一定时间的保温后水冷到室温，制备成最终的成品带钢。
步骤1)所述的锻坯重新加热的温度范围为1100-1200℃，保温时间为3-5h，热轧的开轧温度为1050-1150℃，终轧温度为850-900℃；热轧采用4辊可逆轧机进行7道次的往复 轧制，前两道次压下率为30-50％，后五道次的压下率为20-30％，之后水冷到650-750℃后放入石棉保温8-10h，从而模拟卷曲过程，热轧带钢的厚度为4-5.5mm。
步骤2)所述的冷轧采用四辊轧机进行单向轧制，轧制道次为10-15道次，其中包含3-5道次平整轧制，最终的冷轧带钢厚度为1.0-1.5mm。
步骤3)所述的冷轧带钢奥氏体化温度为870-920℃，奥氏体化保温时间为5-15min。
步骤4)所述氧化铁皮和脱碳层的去除厚度为上下底面各50-100μm，后将预淬火带钢重新加热的退火温度为780-830℃，退火保温时间为3-8min。随后进行盐浴冷却，盐浴冷却速度为100-200℃/s，盐浴保温温度为320-400℃，保温时间为5-10min。
该发明钢的主要特征在于采用添加较多的微合金元素Nb、Ti细化晶粒，来获得高延伸率(A％≥24％)和高强度(≥980MPa)。与传统TRIP钢生产工艺相比，该发明采用的是对冷轧带钢进行两次热处理的方法：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，随后进行表面除磷和去除脱碳层，再重新进行一次加热退火，最终得到成品带钢。该方法存在微合金元素添加量高和两次退火导致的制造成本增加和制造工序难度增加等问题。
中国专利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。
该专利的不足主要在于：
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，是均热时间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，起不到应有的作用，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”，该高强度双相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和其他不可避免杂质。该双相钢的力学性能：屈服强度Rp 0.2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的加热速率再加热至单相奥氏体区850～950℃；之后，钢板在保温不超过5s后立即水冷到室温，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，所以这给热处理本工艺、热处理前工序制造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很高的实际应用价值或实际应用价值不大。
中国专利CN1081931138B公开了“一种980MPa级汽车用冷轧高强Q&P钢及其生产方法”，该钢化学成分质量百分比为：C：0.18～0.24％，Si：0.6～1.3％，Mn：1.6～2.4％，P：0.02～0.04％，S≤0.005％，Nb：0.04～0.07％，N≤0.006％，Als：0.5～1.0％，余量为Fe和其它不可避免的杂质元素。其热轧工序终轧温度870～910℃，卷取温度660～710℃；所述冷轧工序的冷轧压下率≥45％；所述连续退火工序的均热段保温温度770～840℃，过时效段保温温度300～440℃，均热段保温时间为60～225s，过时效段保温时间为300～1225s；所述平整工序平整延伸率为0.3～0.9％。所得到的钢板的屈服强度大于550MPa，抗拉强度大于980MPa，断后伸长率大于18％。
该发明钢的主要特征在于通过传统Q&P工艺，获得良好强塑性配合的结果。由于采 用传统热处理工艺方法，其均热时间和配分时间均很长，同时其合金含量相对较高，这也会增加制造成本、降低制造灵活性。
中国专利申请CN109136779A公开了“一种马氏体基体1100MPa级稀土Q&P钢制备方法”，该发明钢化学成分质量百分比为：C：0.15～0.22％，Si：0.6～1.7％，Mn：1.1～2.4％，Mo：0.1～0.5％，Al：0.1～0.5％，V：0.05～0.11％，Y：0.01～0.05％，P：0.02～0.04％，S≤0.005％，Nb：0.04～0.07％，N≤0.006％，B：0.001～0.006％，余量为Fe和其它不可避免的杂质元素。所得钢板的抗拉强度在1100MPa左右，断后伸长率20％左右。
该发明钢的主要特征在于通过同时添加稀土Y和Mo、V、Nb等合金元素用于细化晶粒，减少Mn元素的含量提升焊接性能，其制造过程需要进行两次铸造。冶炼过程：按照该发明所给出的成分配方，配料后经过转炉冶炼，真空炉二次精炼、铸造得到铸坯；微量元素熔入过程：在电弧熔炼炉中加入微量合金元素粉末(Mo、Al、V、Y、Nb、N，B等)，得到二次铸坯。热轧过程：使用加热炉将铸坯加热到1100-1150℃并进行保温1-3h，随后进行热轧，终轧温度为820-880℃，卷取温度550-650℃，所获得的钢板厚度为1.5-3.0mm，后水淬到室温；冷轧过程：酸洗后进行多道次冷轧，获得厚度为1.2-1.5mm的钢板；
其整个退火过程分为三次：
第一次：两相区锰配分过程：将材料以10-30℃/s加热到AC3和AC1(双相区)之间某一温度并保温3-15min后水淬至室温；
第二次：一次碳配分过程：将材料在MS与Mf之间某一温度T0进行保温10-300s，后将材料水淬至室温；
第三次：二次碳配分过程：将材料在MS与Mf之间某一温度T1(T1温度比T0稍低)进行10-300s的二次碳分配，后将材料水淬至室温。
该发明制造工艺复杂，能耗高，合金含量高且复杂，且多次水淬处理，涉及到材料表面氧化层的去除，带来环境和能耗等方面的诸多问题，导致制造成本增加和制造灵活性的降低。
中国专利申请CN108431248A公开了“一种用于制造具有改善延展性和可成形性的高强度钢板的方法和所获得的钢板”，该发明钢化学成分质量百分比为：C：0.15～0.23％，Mn：2.0～2.8％，Si：1.0～2.1％，Al：0.02～1.0％，Al+Si：1.0～2.1％，Nb：0～0.035％，Mo≤0.3％，Cr≤0.04％，余量为Fe和其它不可避免的杂质元素。所述钢板在退火温度TA下退火以获得包含至少65％的奥氏体和最多至35％的铁素体的组织。使所述钢板以至少20℃/秒的冷却速率从至少600℃的温度淬火至Ms-170℃到Ms-80℃的淬火温度QT，将所述钢板加热至350℃至450℃的配分温度PT，并将所述钢板保持在该配分温度下80-440s 的配分时间Pt，随后立刻将所述钢板冷却至室温。所得钢板的抗拉强度大于1180MPa，断后伸长率大于12％。
该发明钢的主要特征在于采用高Mn、高Si和高Al成分通过传统Q&P工艺，控制最终组织中各相的比例，达到获得良好强塑性配合的结果。由于采用传统热处理工艺方法，其均热时间和配分时间均很长，这会增加制造成本、降低制造灵活性。
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。该发明钢的主要制造步骤如下：
1)将一定化学成分的铸坯锻造成锻坯，重新加热后热轧，水冷后卷曲，得到热轧带钢
2)将所述热轧带钢经酸洗后冷轧成冷轧带钢；
3)将所述冷轧带钢完全奥氏体化后保温一段时间，然后水冷到室温，生成全马氏体组织的预淬火带钢；
4)将所述预淬火带钢表面除鳞，从而去除氧化铁皮层和脱碳层，之后重新加热退火保温一段时间，然后用盐浴冷却到一定温度，经过一定时间的保温后水冷到室温，制备成最终的成品带钢。
步骤1)所述的锻坯重新加热的温度范围为1100-1200℃，保温时间为3-5h，热轧的开轧温度为1050-1150℃，终轧温度为850-900℃；热轧采用4辊可逆轧机进行7道次的往复轧制，前两道次压下率为30-50％，后五道次的压下率为20-30％，之后水冷到650-750℃后放入石棉保温8-10h，从而模拟卷曲过程，热轧带钢的厚度为4-5.5mm。
步骤2)所述的冷轧采用四辊轧机进行单向轧制，轧制道次为10-15道次，其中包含3-5道次平整轧制，最终的冷轧带钢厚度为1.0-1.5mm。
步骤3)所述的冷轧带钢奥氏体化温度为870-920℃，奥氏体化保温时间为5-15min。
步骤4)所述氧化铁皮和脱碳层的去除厚度为上下底面各50-100μm，后将预淬火带钢重新加热的退火温度为780-830℃，退火保温时间为3-8min。随后进行盐浴冷却，盐浴冷却速度为100-200℃/s，盐浴保温温度为320-400℃，保温时间为5-10min。
该发明钢的主要特征在于采用添加较多的微合金元素Nb、Ti细化晶粒，来获得高延伸率(A％≥24％)和高强度(≥980MPa)。与传统TRIP钢生产工艺相比，该发明采用的是对冷轧带钢进行两次热处理的方法：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，随后进行表面除磷和去除脱碳层，再重新进行一 次加热退火，最终得到成品带钢。该方法存在微合金元素添加量高和两次退火导致的制造成本增加和制造工序难度增加等问题。
当前受传统连续退火炉生产线设备能力所限，冷轧Q&P钢产品及退火工艺相关研究都是基于现有工业装备的加热速率(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辊子数目较多等，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，其高温炉段内的辊子数目一般在20～40根不等，使带钢表面质量控制难度增大。
发明内容
本发明的目的在于提供一种抗拉强度≥1180MPa的低碳低合金Q&P钢、抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢以及它们的快速热处理制造方法。本发明通过快速热处理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(包括再结晶形核率和奥氏体相变形核率)，缩短晶粒长大时间，细化晶粒，提高残余奥氏体含量，从而进一步提高材料的强度和塑性。本发明低碳低合金Q&P钢的基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1～3μm，马氏体强化相晶粒周围存在均匀分布的残余奥氏体相和铁素体相，其中片层状按体积分数占比为马氏体组织75～90％、残余奥氏体组织10～25％、铁素体组织3～10％。本发明低碳低合金Q&P钢的屈服强度≥660MPa，抗拉强度为≥1180MPa，延伸率≥18％，强塑积≥24GPa％，具有良好的强韧性匹配以及成型和焊接等用户使用性能。本发明采用快速热处理工艺提高了生产效率，降低同级别钢中的合金含量，从而降低生产成本及热处理前工序制造难度，显著减少炉辊数量，提高材料表面质量。
为达到上述目的，本发明的技术方案是：
抗拉强度≥1180MPa的低碳低合金Q&P钢，其化学成分质量百分比为：C：0.16～0.23％，Si：1.1～2.0％，Mn：1.6～3.0％，P≤0.015％，S≤0.005％，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，该抗拉强度≥1180MPa的低碳低合金Q&P钢的金相组织为马氏体75～90％、残余奥氏体10～25％、铁素体3～10％的多相组织，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，该Q&P钢的金相组织中奥氏体具有良好的热稳定性，-50℃奥氏体转变率低于8％，-190℃奥氏体转变率低 于30％。优选地，该Q&P钢的屈服强度668～1112MPa，抗拉强度1181～1350MPa，延伸率18.9～24.2％，强塑积24.1～28.97GPa％。
优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，C的含量范围选自0.17～0.23％、0.19～0.21％和0.18～0.21％。优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Si的含量范围选自1.1～1.7％、1.3～1.5％、1.4～2.0％和1.6～1.8％。优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Mn的含量范围选自1.6～2.2％、1.8～2.0％、2.4～3.0％和2.6～2.8％。
优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Cr的含量≤0.35％，如≤0.25％；Mo的含量≤0.25％；Nb的含量≤0.06％，如≤0.04％；Ti的含量≤0.065％，如≤0.04％，如0.006～0.016％；V的含量≤0.055％，如≤0.035％。
优选地，本发明所述的抗拉强度≥1180MPa的低碳低合金Q&P钢通过下述工艺获得：
1)冶炼、铸造
按上述化学成分冶炼并铸造成板坯；
2)热轧、卷取
卷取温度550～680℃；
3)冷轧
冷轧压下率为40～85％；
4)快速热处理
冷轧后的钢板快速加热至770～845℃，所述快速加热采用一段式或两段式；采用一段式快速加热时加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以50～500℃/s的加热速率从550～625℃加热至770～845℃；之后进行均热，均热温度为770～845℃，均热时间为10～60s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃，随后以50～200℃/s的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s，随后以10～30℃/s的加热速率加热至300～470℃进行回火处理，回火时间10～60s；回火结束后以30～100℃/s的冷却速率冷却至室温。
优选的，步骤2)中，所述热轧终轧温度≥Ar3。
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)所述的快速热处理全过程用时为71～186s。
优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～625℃；第二段以50～300℃/s的加热速率从550～625℃加热至770～845℃。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以30～300℃/s的加热速率从室温加热至550～625℃；第二段以80～300℃/s的加热速率从550～625℃加热至770～845℃。
优选的，步骤4)中，所述钢板快速冷却速率为50～150℃/s。
在一些实施方案中，本发明的抗拉强度≥1180MPa的低碳低合金Q&P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≤0.005％，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。优选的，此低碳低合金Q&P钢中，C含量为0.19～0.21％。优选的，此低碳低合金Q&P钢中，Si含量为1.3～1.5％。优选的，此低碳低合金Q&P钢中，Mn含量为1.8～2.0％。优选地，此Q&P钢的金相组织为马氏体75～85％、残余奥氏体10～25％和铁素体3～10％的多相组织，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，此Q&P钢的金相组织中奥氏体具有良好的热稳定性，-50℃奥氏体转变率低于8％，-190℃奥氏体转变率低于30％。优选地，此Q&P钢的屈服强度为668～1002MPa，抗拉强度为1181～1296MPa，延伸率为18.9～24.2％，强塑积为24.1～28.6GPa％。
在一些实施方案中，本发明的抗拉强度≥1180MPa的低碳低合金Q&P钢的化学成分质量百分比为：C：0.16～0.23％，Si：1.4～2.0％，Mn：2.4～3.0％，Ti：0.006～0.016％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Nb、V中的一种或两种，且，Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，此抗拉强度≥1180MPa的低碳低合金Q&P钢为抗拉强度≥1280MPa的低碳低合金Q&P钢。优选的，此低碳低合金Q&P钢中，C含量为0.18～0.21％。优选的，此低碳低合金Q&P钢中，Si含量为1.6～1.8％。优选的，此低碳低合金Q&P钢中，Mn含量为2.6～2.8％。此Q&P钢的金相组织为马氏体80～90％、残余奥氏体10～20％、铁素体3～5％的多相组织，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，此Q&P钢的金相组织中奥氏体具有良好的热稳定性，-50℃奥氏体转变率低于8％，-190℃奥氏体转变率低于30％。优选地，此Q&P钢的屈服强度754～1112MPa，抗拉强度1281～1350MPa， 延伸率19～22.2％，强塑积24.8～28.97GPa％。
本发明另一方面提供抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢，其化学成分质量百分比为：C：0.16～0.23％，Si：1.1～2.0％，Mn：1.6～3.0％，P≤0.015％，S≤0.005％、优选≤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，所述热镀锌Q&P钢的金相组织为马氏体、铁素体和奥氏体三相组织，其基体组织分布均匀，出现片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，所述热镀锌Q&P钢的金相组织按体积分数占比为马氏体45～75％、铁素体15～30％、奥氏体10～25％的三相组织。优选地，所述热镀锌Q&P钢的屈服强度≥720MPa，抗拉强度≥1180MPa，延伸率≥19％，强塑积≥23.0GPa％。优选地，所述热镀锌Q&P钢的屈服强度为721～956MPa，抗拉强度为1184～1352MPa，延伸率为19～22.5％，强塑积为23.6～28.9GPa％。优选地，所述热镀锌Q&P钢金相组织中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≥1180MPa的低碳低合金热镀锌Q&P钢中，C的含量范围选自0.17～0.23％、0.19～0.21％和0.18～0.21％。优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Si的含量范围选自1.1～1.7％、1.3～1.5％、1.4～2.0％和1.6～1.8％。优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Mn的含量范围选自1.6～2.2％、1.8～2.0％、2.4～3.0％和2.6～2.8％。
优选地，所述抗拉强度≥1180MPa的低碳低合金Q&P钢中，Cr的含量≤0.35％，如≤0.25％；Mo的含量≤0.25％；Nb的含量≤0.06％，如≤0.04％；Ti的含量≤0.065％，如≤0.04％，如0.006～0.016％；V的含量≤0.055％，如≤0.035％。
在一些实施方案中，所述抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢通过以下方法制备得到：
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≥Ar3，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770～845℃，所述快速加热采用一段式或两段式；采用一段 式快速加热时，加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以30～500℃/s(如50～500℃/s)的加热速率从550～625℃加热至770～845℃；之后进行均热，均热温度：770～845℃，均热时间：10～60s；
均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃，随后以50～200℃/s的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s，随后以10～30℃/s的加热速率加热至460～470℃进行配分处理，配分时间10～60s；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10～300℃/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，所述快速热处理、热镀锌全过程用时为43～186s。
优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以15～300℃/s的加热速率从室温加热至550～625℃，第二段以50～300℃/s的加热速率从550～625℃加热至770～845℃。
优选的，步骤4)中，所述快速加热采用两段式加热：第一段以30～300℃/s的加热速率从室温加热至550～625℃，第二段以80～300℃/s的加热速率从550～625℃加热至770～845℃。
优选的，步骤4)中，所述带钢或钢板快速冷却阶段冷却速率为50～150℃/s。
在一些实施方案中，所述抗拉强度≥1180MPa的低碳低合金热镀锌Q&P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≤0.005％，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。优选的，该热镀锌Q&P钢中，C含量为0.19～0.21％。优选的，该热镀锌Q&P钢中，Si含量为1.3～1.5％。优选的，该热镀锌Q&P钢中，Mn含量为1.8～2.0％。优选地，该热镀锌Q&P钢的金相组织中按体积分数占比为马氏体45～75％、铁素体15～30％和奥氏体10～25％的三相组织，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，该热镀锌Q&P钢的屈服强度为721～805MPa，抗拉强度为1184～1297MPa，延伸率为19.1～22.4％，强塑积为 23.6～28GPa％。优选地，该热镀锌Q&P钢金相组织中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％。
在一些实施方案中，所述抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢的化学成分质量百分比为：C：0.16～0.23％，Si：1.4～2.0％，Mn：2.4～3.0％，Ti 0.006～0.016％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、Nb、V中的一种或两种，且Cr+Mo+Ti+Nb+V≤0.5％，余量为Fe和其它不可避免的杂质。优选地，该抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢为抗拉强度≥1280MPa的低碳低合金热镀锌Q&P钢。优选的，该热镀锌Q&P钢中，C含量为0.18～0.21％。优选的，该热镀锌Q&P钢中，Si含量为1.6～1.8％。优选的，该热镀锌Q&P钢中，Mn含量为2.6～2.8％。优选地，该热镀锌Q&P钢的金相组织为马氏体、铁素体和奥氏体三相组织(马氏体组织占比75～90％、残余奥氏体组织占比10～25％、铁素体组织占比3～10％)，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主。优选地，该热镀锌Q&P钢的屈服强度为802～956MPa，抗拉强度为1280～1352MPa，延伸率19～22.5％，强塑积25.2～28.9GPa％。优选地，该热镀锌Q&P钢金相组织中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％。
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。碳在钢中的相有两种存在方式：铁素体和渗碳体。碳含量对钢的力学性能影响十分大，随着含碳量的升高，马氏体和珠光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，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Q&P钢而言，碳元素是马氏体基体最有效的强化元素之一，它固溶于奥氏体中，扩大奥氏体相区，极大的提高奥氏体稳定性，使珠光体和贝氏体的转变C曲线右移，推迟珠光体和贝氏体的转变，降低Ms点温度。含碳量太低会使残余奥氏体的稳定性降低，含碳量过高会使马氏体中出现孪晶，降低钢的塑性、韧性和焊接性。综合考虑将含碳量限定在0.16～0.23％范围之内。在一些实施方案中，C的含量为0.18～0.21％。在另外一些实施方案中，C的含量为0.19～0.21％。
Mn：锰可以与铁形成固溶体，进而提高碳钢中铁素体与奥氏体的强度及硬度，并使 钢材在热轧之后的冷却过程中获得较细小且强度较高的珠光体，而且珠光体的含量也会随着Mn含量的增加而有所增加。锰同时又是碳化物的形成元素，锰的碳化物能够溶入渗碳体，从而间接地增强马氏体和珠光体等强化相的强度。锰还可以强烈增强钢的淬透性，进一步提高其强度。在一些实施方案中，Mn的含量为1.8～2.0％。在另外一些实施方案中，C的含量为2.6～2.8％。
对于Q&P钢而言，添加锰元素可降低马氏体转变温度Ms，增加残余奥氏体的含量，提高残余奥氏体的稳定性，且锰元素对钢的韧性影响不大。但锰含量较高时，有使钢材晶粒粗化的趋势，并且增加钢的过热敏感性，当熔炼浇注与热轧之后冷却不当时，容易使碳钢中产生白点。本发明将锰含量限定在1.6～3.0％范围之内。
Si：硅在铁素体或奥氏体中形成固溶体，从而增强钢的屈服强度与抗拉强度，而且硅可增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，但是当钢中硅的含量过高时，其塑性与韧性会显著下降。对于Q&P钢而言：
第一，硅元素是非碳化物形成元素，在碳化物中的溶解度极低，在QP钢等温过程中，能够抑制Fe3C的形成，使未转变的奥氏体富碳，从而大大提高奥氏体的稳定性，使其能够在室温保留下来；
第二，硅元素是铁素体形成元素，可以提高残余奥氏体的稳定性，起到固溶强化的作用，从而提高钢的强度；
第三，硅元素有缩小奥氏体相区，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坚硬的氧化层、较差的表面性能、降低热轧钢板的润湿性和表面质量。硅对奥氏体长大速率没有明显影响，但对奥氏体的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；本发明将硅含量限定在1.1～2.0％范围之内。在一些实施方案中，Si的含量为1.3～1.5％。在另外一些实施方案中，Si的含量为1.6～1.8％。
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。加入其他合金元素时，铬元素提高钢的强度和硬度效果较显著。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。在一些是实施方案，Cr的含量≤0.35％，如≤0.25％。
Mo：钼元素能抑制铁的自扩散和其他元素的扩散速度。Mo原子半径比α-Fe原子大，当Mo溶解在α固溶体时，使固溶体发生强烈的晶格畸变，同时Mo能增加晶格原子键引力，提高α铁素体的再结晶温度。Mo在珠光体型、铁素体型、马氏体型钢中，甚至在高合金奥氏体钢中的强化作用也十分明显。Mo在钢中的良好作用还需视与钢中其他合金元素间的相互作用而定。在钢中加入强碳化物形成元素V、Nb、Ti时，Mo的固溶强化作用更加显著。这是因为当强碳化物形成元素与C结合成稳定的碳化物时，能促进Mo更有效地溶入固溶体中，从而更有利于钢的热强性提高。加入Mo还可以增加钢的淬透性，但效果没有C和Cr显著。Mo会抑制珠光体区的转变，使中温区转变加快，因而含Mo钢在冷却速度较大的情况下也能形成一定数量的贝氏体，并且消除铁素体的形成，这是Mo对低合金耐热钢热强性产生有利影响的原因之一。Mo还能显著降低钢的热脆倾向，并减小珠光体球化速度。当Mo含量在0.15％以下时，对钢的焊接性能无不利的影响。本发明中钼元素为有益且非必要添加元素，考虑成本增加等因素添加量不宜过多。在一些实施方案中，Mo的含量≤0.25％
Nb：Nb元素是碳化物和氮化物的形成元素，且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。加入Nb可以阻止奥氏体晶粒长大，提高钢材晶粒的粗化温度。Nb元素与碳生成十分稳定的NbC，在钢中添加微量的Nb元素可以利用其析出强化的效果，提高基体的强度。Nb元素对铁素体再结晶的长大和奥氏体的晶粒长大有明显的阻碍作用，能够细化晶粒，提高钢的强度和韧性；Nb元素可以影响晶界的移动性，对相变行为和碳化物的形成也有影响。Nb可使碳在残余奥氏体中的含量升高，阻碍贝氏体的形成，促使马氏体形核，获得弥散分布的马氏体组织，并且能够提高残余奥氏体的稳定性，通过添加Nb元素来提高双相钢的强度可以做在较低含量的马氏体和低C含量的条件下得到一定强度的双相钢，提高双相钢的强韧性；同时添加Nb元素的另外一个好处是可以在一个较宽的退火温度范围内提高钢的强度。本发明中Nb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≤0.06％，如≤0.04％。
Ti：Ti是微合金元素，属于封闭γ区的铁素体形成元素，它可提高钢的临界点，钢中的Ti和C可形成十分稳定的TiC，在一般热处理的奥氏体化温度范围内，TiC极难溶解。由于TiC颗粒使奥氏体晶粒细化，奥氏体分解转变时，新相晶核形成的机会增多，这些都加速了奥氏体转变。另外，Ti可与C，N形成TiC、TiN析出相，比Nb、V的碳氮化物更稳定，显著降低C在奥氏体中的扩散速度，使奥氏体形成速度大幅度降低，形成的碳氮化物在基体中沉淀，钉扎在奥氏体的晶界，阻碍奥氏体晶粒长大。在冷却过程中，析出的TiC具有沉淀强化作用；在回火过程中，Ti减缓C在α相中的扩散，减缓Fe、Mn等碳化物的析出与长大，增加回火稳定性，并可通过析出TiC而起到二次硬化作用。通过Ti的微合金化可提高钢的高温强度。在钢中添加微量的Ti，一方面，可在减少碳当量含量的同时提高强度、提高钢的焊接性能；另一方面，将不纯物质如氧、氮、硫等固定起来，从而改善钢的可焊性；其次，由于其微观质点的作用，例如TiN在高温下的未溶解性，可阻止热影响区晶粒的粗化，提高热影响区的韧性，从而改善钢的焊接性能。在一些实施方案中，Ti的含量≤0.065％，如≤0.04％。在一些实施方案中，Ti添加时，其添加量可在0.006～0.016％的范围内。
微合金元素V：V是铁素体稳定元素，且是强碳化物形成元素，具有强烈的细化晶粒作用，可使钢的组织致密。钢中添加V可使钢的强度、塑性和韧性同时得到改善。钒还可以提高结构钢的高温强度。钒不能提高淬透性。在钢中添加微量的微合金元素V，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V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等使用性能。在钢中添加微量的V，一方面，可在减少碳当量含量的同时提高强度、提高钢的焊接性能；另一方面，将不纯物质如氧、氮、硫等固定起来，从而改善钢的可焊性；其次，由于其微观质点的作用，例如V(CN)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，V的含量≤0.055％，如≤0.035％。
在钢中添加微量的微合金元素Nb、V、Ti，可保证钢在碳当量较低的情况下，通过其碳、氮化物质点(尺寸小于5nm)的弥散析出及Nb、V、Ti的固溶，细化晶粒，极大地提高钢的强度、韧性，特别是低温韧性，使钢具有良好的可焊性、使用性。Nb、V、Ti是碳化物和氮化物的形成元素，这些元素在比较低的浓度下就能满足这种要求Nb、V、Ti为强碳化物形成元素，常温时，在钢中大部分以碳化物、氮化物、碳氮化物形式存在，少部分 固溶在铁素体中。对于Q&P钢而言，添加这些微合金化元素，能够通过晶粒细化和沉淀强化铁素体基体。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动力。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。
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)来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、弥散分布的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细化。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CN107794357B和美国专利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，水淬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Q&P钢的主要相组织为马氏体(体积分数占比75～90％)和残余奥氏体(体积分数占比10～25％)，同时含有极少量的铁素体(体积分数占比3～10％)，因此严格说来其相组织为多相组织，其基体组织分布均匀，出现明显片层状回火马氏体，晶粒粒径为1～3μm，马氏体强化相晶粒周围由均匀分布的铁素体相包围，马氏体强化相晶粒以片状组织结构为主。
本发明所述的抗拉强度≥1180MPa的低碳低合金Q&P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
卷取温度550～680℃；
3)冷轧
冷轧压下率40～85％，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770～845℃，所述快速加热采用一段式或两段式；采用一段式快速加热时加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以50～500℃/s的加热速率从550～625℃加热至770～845℃；
b)均热
在奥氏体和铁素体两相区目标温度770～845℃进行均热，均热时间为10～60s；
c)冷却
带钢或钢板均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃(如720～770℃)，随后以50～200℃/s(如50～150℃/s)的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s；
d)回火
保温结束后，将带钢或钢板以10～30℃/s的加热速率加热至300～470℃进行回火处理，回火时间10～60s
e)回火结束后带钢或钢板冷却至室温，冷却速率30～100℃/s。
所述低碳低合金Q&P钢的制造方法中，优选的，步骤2)中，所述热轧终轧温度≥Ar3。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，所述快速热处理全过程用时为71～186s。优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。优选的，步骤4)中，所述快速加热采用两段式加热，第一段以15～300℃/s的加热速率从室温加热至550～625℃；第二段以50～300℃/s的加热速率从550～625℃加热至770～845℃。优选的，步骤4)中，所述快速加热采用两段式加热，第一段以30～300℃/s的加热速率从室温加热至550～625℃；第二段以80～300℃/s的加热速率从550～625℃加热至770～845℃。优选的， 步骤4)中，所述快速加热最终温度为790～845℃。优选的，步骤4)中，所述带钢或钢板快速冷却速率为50～150℃/s。优选的，步骤4)所述的均热过程中，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)所述的均热过程中，带钢或钢板在均热时间段内进行小幅度升温或小幅度降温，升温后温度不超过845℃，降温后温度不低于770℃。优选的，所述均热时间为10～40s。
本发明所述的抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢的快速热处理、热镀锌制造方法，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≥Ar3，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770～845℃奥氏体和铁素体两相区目标温度，所述快速加热采用一段式或两段式；
采用一段式快速加热时，加热速率为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770～845℃；
b)均热
在奥氏体和铁素体两相区目标温度770～845℃进行均热，均热时间为10～60s；
c)冷却
带钢或钢板均热结束后以5～15℃/s冷却速率缓冷至720～770℃；随后以50～200℃/s(如50～150℃/s)冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s(如2～8s)；
d)配分
保温结束后，将带钢或钢板以10～30℃/s的加热速率加热至460～470℃进行配分处理，配分时间10～60s；
e)热镀锌
配分结束后将带钢或钢板浸入锌锅进行热镀锌；
f)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10～300℃/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的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
该低碳低合金热镀锌Q&P钢的快速热处理、热镀锌制造方法中，优选的，所述快速热处理、热镀锌全过程用时为43～186s。优选的，步骤2)中，所述卷取温度为580～650℃。优选的，步骤3)中，所述冷轧压下率为60～80％。优选的，步骤4)中，所述快速加热采用一段式加热时加热速率为50～300℃/s。优选的，步骤4)中，所述快速加热采用两段式加热，第一段以15～300℃/s的加热速率从室温加热至550～625℃，第二段以50～300℃/s的加热速率从550～625℃加热至770～845℃。优选的，步骤4)中，所述快速加热采用两段式加热，第一段以30～300℃/s的加热速率从室温加热至550～625℃，第二段以80～300℃/s的加热速率从550～625℃加热至770～845℃。优选的，步骤4)中，所述快速加热最终温度为790～845℃。优选的，步骤4)中，所述快速冷却阶段冷却速率为50～150℃/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，带钢或钢板均热时间段内进行小幅度升温或小幅度降温，升温后温度不超过845℃，降温后温度不低于770℃。优选的，所述均热时间为10～40s。优选的，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≥1180MPa的低碳低合金Q&P钢及低碳低合金热镀锌Q&P钢的制造方法中：
1、加热速度控制
连续加热过程的再结晶动力学可以由受加热速率影响的关系式来定量描述，连续加热过程中铁素体再结晶体积分数与温度T的函数关系式为：
Figure PCTCN2022084518-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。
传统慢速加热条件下，变形基体都先回复、再结晶及晶粒长大，而后发生铁素体向奥氏体的相转变，且相变形核点主要集中在已经长大的铁素体晶界处，形核率较低，因此最终得到的晶粒组织比较粗大。
快速加热条件下，变形基体还没有充分回复就开始再结晶，再结晶还没有完成或晶粒长大还没开始，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时晶粒细小、晶界面积大，因此形核率显著提高，晶粒明显细化。特别是铁素体再结晶与奥氏体相变过程发生重叠后，由于铁素体晶体内保留了大量位错等晶体缺陷，为奥氏体提供了大量的形核点，使得奥氏体呈现爆发式形核，奥氏体晶粒进一步细化。同时保留下来的高密度位错线缺陷也成为碳原子高速扩散的通道，使得每一个奥氏体晶粒都能快速生成并长大，因此增大奥氏体体积分数。
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50～300℃/s；加热温度从500～625℃到奥氏体化温度770～845℃，加热速率对再结晶形核、相变形核及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。
均热温度取决于C含量，传统工艺中一般将均热温度设置在A C3以上30～50℃，本发明利用快速加热技术在铁素体中形成大量位错，为奥氏体转变提供了形核功，所以只需要将温度加热到A C1到A C3之间。本发明Q&P钢中C含量为0.16～0.23％，A C1和A C3分别是730℃和870℃左右。Q&P钢中有大量未溶解的细小均匀分布的碳化物，在均热处理过程中，能够对奥氏体晶粒的长大起到机械阻碍的作用，有利于细化合金钢的晶粒度，但是如果均热温度过高，就会使未溶解的碳化物数目大量减少，削弱这种阻碍作用，增强晶粒 的长大倾向，进而降低钢的强度。当未溶碳化物的数量过大时，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，达到改善合金钢的强度与韧性等力学性能的目的。
均热温度的选取还应以获得细小均匀的奥氏体晶粒为目的，以达到在冷却之后能够得到较高体积分数且均匀细小的马氏体组织的目的。过高的均热温度会使奥氏体晶粒粗大，淬火过程中工件容易开裂，淬火后获得的马氏体组织也会较粗大，使钢的力学性能不佳；还会降低残余奥氏体的数量，降低工件的硬度与耐磨性。过低的均热温度，又会使奥氏体溶入的碳以及合金元素含量不足，令奥氏体中的合金元素浓度分布不均，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Ac 3+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770～845℃作为均热温度，以期获得合理的淬火工艺及理想组织性能。
3、均热时间控制
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；均热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；但是如果均热时间过短，会使钢中合金元素分布不均，导致奥氏体化不充分；均热时间过长又容易导致奥氏体晶粒粗大。均热时间的长短也与钢中碳以及合金元素的含量有关，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。所以均热时间的控制需严格结合均热温度、快速冷却及快速加热过程综合考虑制定，才能最终获得理想的组织和元素分布。综上，本发明将均热保温时间定为10～60s。
4、快速冷却速度控制
为了获得马氏体，快冷时试样的冷速必须大于临界冷却速度才能够得到马氏体组织，临界冷却速度主要取决于材料成分，本发明中的Si含量为1.1～2.0％，Mn含量为1.6～3.0％，含量相对较高，所以Si和Mn很大程度加强了Q&P钢的淬透性，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。冷却速率太低无法获得马氏体组织，会导致强度下降，力学性能无法满足要求；而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)引起板形严重不良，冷却不均匀时板形不良尤其严重，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50～200℃/s。
5、回火温度控制
通常合金钢在150℃以下进行回火时，由于温度过低，合金元素无法进行扩散，只有碳元素还具有一定的扩散能力，因此低温回火钢虽然具备较高的硬度，但其脆性过大，韧性很差，无法满足工件的使用性能要求。当在200℃以上温度进行回火时，马氏体含有的碳元素与其他合金元素会开始大量析出，使残余应力减小直至消失，回火钢的硬度也会随回火温度的上升而逐渐下降，但韧性增强。而当回火温度达到500℃左右时，马氏体分解结束，渗碳体逐渐聚集长大，α相开始发生回复过程，继续升高温度，α相开始再结晶，形成多边形铁素体，强度显著下降。回火温度越高，α相与渗碳体相越粗大，回火钢的硬度也会越低，本发明最终目的是同时获得较好的强度和塑性，所以本发明将回火温度设置在300～470℃。
6、回火时间控制
钢在回火过程中，回火时间起到三方面的作用：(1)保证组织转变进行充分；(2)降低或消除内应力；(3)与回火温度配合使工件获得所需要的性能。本发明钢中由于采用快速加热技术使得奥氏体晶粒细化，从而将一次快速冷却后生成的残余奥氏体与马氏体的间距缩短，碳原子由过饱和马氏体向残余奥氏体扩散配分的效率提高，因此回火过程所需时间也大大减小。但如果回火时间过短难以消除内应力、降低工件的脆硬性，综合考虑，本发明将回火时间设置在10～60s。
7、配分温度控制
通常较高合金含量的Q&P钢在150℃以下进行碳分配(回火)，由于温度过低，合金元素无法进行扩散，只有碳元素还具有一定的扩散能力，因此低温配分钢虽然具备较高的硬度，但其脆性过大，韧性也很差，无法满足工件的使用性能要求。当在200℃以上温度进行配分时，马氏体含有的碳元素与其他合金元素会开始大量析出，使残余应力减小直至消失，配分钢的硬度也会随配分温度的上升而逐渐下降。当配分温度达到500℃左右时，马氏体分解结束，渗碳体逐渐聚集长大，α相开始发生回复过程，继续升高温度，α相开始再结晶，形成多边形铁素体。配分温度越高，α相与渗碳体相越粗大，配分钢的硬度也会越低，本发明的配分工艺最主要的目的是为了将已经获得的马氏体中的碳扩散到尚未发生马氏体转变的残余奥氏体中，使得马氏体中的碳降低塑性提高，同时扩散到残余奥氏体中的碳浓度提高，增强其稳定性，使得最终产品同时获得较好的强度和塑性，即良好的强塑性配合，再结合热镀锌温度，所以将配分温度设置在460～470℃。
8、配分时间控制
钢在配分过程中，配分时间起到三方面的作用：(1)保证组织转变进行充分；(2)尽量降低或消除内应力；(3)与配分温度配合使工件获得所需要的性能。本发明中由于采用快速加热技术使得奥氏体晶粒细化，从而将一次快速冷却后生成的残余奥氏体与马氏体的间距缩短，碳原子由过饱和马氏体向残余奥氏体扩散配分的效率提高，因此配分过程所需时间也大大减小。但如果配分时间过短难以消除内应力、降低工件的脆硬性，综合考虑，本发明将配分时间设置在10～60s。
10、热镀锌和合金化控制
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。
通过本发明方法，可降低同级别钢中的合金含量，细化晶粒、获得良好的软、硬相组织构成及强度和韧性的匹配；同时，通过对传统连续热镀锌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)，提高传统连续热镀锌机组的生产效率，降低生产成本及能耗，显著减少连续退火炉炉辊数量，特别是高温炉段炉辊数量，这可以降低能耗和对设备的投入。
同时，通过建立快速热处理、热镀锌工艺技术的新型连续退火热镀锌机组，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热镀锌前工序制造难度，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，通过建立采用快速热处理工艺技术的新型连续退火机组，可实现机组短小精悍、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，进一步提高材料的强度和塑性，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Q&P钢的金相组织为马氏体占75～90％、残余奥氏体占10～25％、铁素体占3～10％的多相组织，所获得的热镀锌Q&P钢的金相组织为细化的马氏体、铁素体和奥氏体三相组织，优选马氏体体积分数占比45～75％、残余奥氏体体积分数占比10～25％、铁素体体积分数占比15～30％。所获得的Q&P钢及热镀锌Q&P钢的基体组织分布均匀，出现明显片层状回火马氏体，且晶粒尺寸细化到1～3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相，马氏体强化相晶粒以片状组织结构为主；组织中的奥氏体具有块状、条状、颗粒状等多种形态，具有良好的热稳定性，-50℃奥氏体转变率低于8％，-190℃奥氏体转变率低于30％，且可在不同应变条件下持续发生TRIP效应，因此产品力学性能及用户使用性能优异。
(2)相比于传统热处理方式所得Q&P钢，本发明得到的Q&P钢合金成分大幅度降低，晶粒尺寸减小40～80％，性能优异；其屈服强度为668～1112MPa，抗拉强度为1181～1350MPa，延伸率为18.9～24.2％，强塑积为24.1～28.97GPa％。相比于传统连续退火热镀锌方式所得的热镀锌Q&P钢，在前工序制造条件不变的前提下，通过本发明快速热处理后得到的Q&P钢的平均晶粒尺寸为1-3μm，平均晶粒尺寸减小10～40％，可获得良好的细晶强化的效果；其屈服强度≥720MPa，抗拉强度≥1180MPa，延伸率≥19％，强塑积≥23.0GPa％；优选地，所述热镀锌Q&P钢的屈服强度为721～956MPa，抗拉强度为1184～1352MPa，延伸率为19～22.5％，强塑积为23.6～28.9GPa％。
(3)本发明所述的抗拉强度≥1180MPa的低碳低合金Q&P钢快速热处理工艺及低碳低合金热镀锌Q&P钢快速热处理工艺热处理全过程用时可分部缩短至71～186s和43～186s，大大降低了整个热处理工艺过程的时间(传统连续退火工艺时间通常在5～8min)，显著提高了生产效率、减少了能耗，降低了生产成本。
(4)相比于传统的Q&P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％，整个热处理工序时间缩短至71～186s；相比于传统的热镀锌Q&P钢及其热处理工艺，本发明的快速热处理方法缩短了连续热镀锌退火炉加热段和均热段的长度和时间(和传统连续热镀锌退火炉相比，加热段和均热段的长度缩短可达60～80％)及整个热处理工序时间。因此，本发明可节能减排降耗，显著降低炉子等设备的一次性投资，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)相比于传统工艺生产的Q&P钢、热镀锌Q&P钢及其热处理工艺，采用快速热 处理工艺技术可以减少加热过程和均热过程时间、缩短炉子长度、减少炉辊数量，使得炉内产生表面缺陷的几率减少，产品表面质量将显著提高；另外由于产品晶粒的细化和材料合金含量的减少，采用本发明技术得到的Q&P钢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。
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。
本发明得到的抗拉强度≥1180MPa的低碳低合金Q&P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Q&P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Q&P钢显微组织图片。
图3是本发明实施例一试验钢K按实施例7所生产的Q&P钢显微组织图片。
图4是本发明实施例一试验钢R按实施例8所生产的Q&P钢显微组织图片。
图5是本发明实施例一试验钢P按实施例22所生产的Q&P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Q&P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Q&P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Q&P钢显微组织图片。
图9是本发明实施例二试验钢K按实施例7所生产的Q&P钢显微组织图片。
图10是本发明实施例二试验钢R按实施例8所生产的Q&P钢显微组织图片。
图11是本发明实施例二试验钢P按实施例22所生产的Q&P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Q&P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Q&P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所生产的热镀纯锌Q&P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双相钢(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所生产的热镀纯锌双相钢(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双相钢(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Q&P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Q&P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双相钢(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所生产的热镀纯锌双相钢(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双相钢(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为本实施例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匹配。通过本发明的方法获得的Q&P钢屈服强度为668～1002MPa，抗拉强度为1181～1296MPa，延伸率为18.9～24.2％，强塑积为24.1～28.6GPa％。
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材料组织存在非常大的差别。经过本发明实施例处理后的A钢其组织(图1)主要为马氏体基体上弥散分布的细小、均 匀的奥氏体组织及少量的碳化物组成，奥氏体、马氏体晶粒组织及碳化物都非常细小且均匀分布于基体中，这对提高材料强度和塑性都是非常有利的。而经过传统工艺处理的A钢组织(图2)则分布相对不均匀，存在少量大块的白色铁素体组织，该铁素体晶界上分布着黑色马氏体和奥氏体组织。采用传统工艺处理的组织特点是：晶粒相对粗大，且存在一定的组织分布不均匀现象。
图3为典型成分K钢经过实施例7获得的组织图，图4为典型成分R钢经过实施例8获得的组织图。图5为典型成分P钢经过实施例22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7、8、22、23均为整个热处理周期较短的工艺。从图中可见，采用本发明方法，经过短时间的快速退火处理可获得更加均匀、细小、弥散分布的各相组织。因此本发明的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 PCTCN2022084518-appb-000002
Figure PCTCN2022084518-appb-000003
Figure PCTCN2022084518-appb-000004
表4
Figure PCTCN2022084518-appb-000005
表5
Figure PCTCN2022084518-appb-000006
实施例二
本发明试验钢的成分参见表6，本发明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9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Q&P钢的屈服强度754～1112MPa，抗拉强度1281～1350MPa，延伸率19～22.2％，强塑积24.8～28.97GPa％。
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材料组织存在非常大的区别。经过本发明实施例处理的获得钢的组织主要为铁素体基体上弥散分布的细小、均匀的马氏体组织及少量的碳化物组成，马氏体晶粒组织及少量碳化物都非常细小且均匀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钢组织则分布相对不均匀，马氏体体晶粒相对较大，马氏体晶界上分布着少量的残余奥氏体和碳化物组织，且分布不均匀。采用传统工艺处理的组织特点是：晶粒相对粗大，且存在一定的组织分布不均匀现象。
图9为典型成分K钢经过实施例7获得的组织图，图10为典型成分R钢经过实施例8获得的组织图。图11为典型成分P钢经过实施例22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7、8、22、23均为整个热处理周期较短的工艺。从图中可见，采用本发明方法，经过短时间的快速退火处理可获得更加均匀、细小、弥散分布的各相组织。因此本发明的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 PCTCN2022084518-appb-000007
Figure PCTCN2022084518-appb-000008
Figure PCTCN2022084518-appb-000009
表9
Figure PCTCN2022084518-appb-000010
表10
Figure PCTCN2022084518-appb-000011
从实施例一和二的结果可知，可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，可以极大的缩短传统连续退火炉加热及均热段的长度，提高传统连续退火机组的生产效率，降低生产成本及能耗，减少连续退火炉的炉辊数量，这可以提高带钢表面质量的控制能力，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，使得该连续热处理机组具有短小精悍、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成本和制造难度，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不超过10根，可明显提高带钢表面质量。同时，在极短时间内所完成的再结晶和奥氏体化过程的快速热处理工艺方法也将提供更加灵活及柔性化的高强钢组织设计方法，进而在无需改变合金成分以及轧制工艺等前工序条件的前提下改善材料组织，提高材料性能。
以Q&P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Q&P钢的开发和生产提供更大的空间。
实施例三
本实施例试验钢的成分参见表11，本发明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QP钢产品的主要性能。
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Q&P钢的屈服强度为721～805MPa，抗拉强度为1184～1297MPa，延伸率为19.1～22.4％，强塑积为23.6～28GPa％。
图13和图14为典型成分A钢经过实施例1和传统工艺例1的组织图。从图上看，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)：基体组织分布均匀，组织出现明显片层状回火马氏体，晶粒粒径为1-3μm。马氏体强化相晶粒周围存在均匀分布的铁素体相。由于部分原奥氏体长大后形成的马氏体稳定性下降使得热处理后组织中出现少量回火马氏体，剩余马氏体强化相仍以片状形貌为主，铁素体、马氏体晶粒组织及碳化物都非常细小且均匀分布于基体中，这对提高材料强度和塑性都是非常有利的。
而经过传统工艺处理的钢组织(图14)则为典型的Q&P钢组织图，板条马氏体晶粒粗大，奥氏体及碳化物沿马氏体晶界分布，多相组织分布不均匀。
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；从图上可以看出，采用本发明方法，可获得非常均匀、细小、弥散分布的各相组织。因此本发明的热镀锌Q&P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，进而改善了材料组织，提高了材料性能。
Figure PCTCN2022084518-appb-000012
Figure PCTCN2022084518-appb-000013
Figure PCTCN2022084518-appb-000014
Figure PCTCN2022084518-appb-000015
表14
Figure PCTCN2022084518-appb-000016
表15
Figure PCTCN2022084518-appb-000017
Figure PCTCN2022084518-appb-000018
实施例四
本实施例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热镀纯锌GI产品的主要性能，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QP钢产品的主要性能。
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Q&P钢屈服强度可达802～956MPa，抗拉强度为1280～1352MPa，最大延为19～22.5％，强塑积25.2～28.9GPa％。
图18和图19为典型成分A钢经过实施例1和传统工艺例1的组织图。从图上看，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)：由马氏体、奥氏体及少量铁素体和碳化物组成，基体组织分布均匀，组织出现明显片层状回火马氏体，晶粒粒径为1-3μm。绝大多数强化相晶粒周围均由铁素体包围。由于部分原奥氏体长大后形成的马氏体稳定性下降使得热处理后组织中出现少量回火马氏体，剩余强化相仍以块状形貌为主，铁素体、马氏体晶粒组织及碳化物都非常细小且均匀分布于基体中，这对提高材料强度和塑性都是非常有利的。
而经过传统工艺处理的钢组织(图19)则为典型的Q&P钢组织图，板条马氏体晶粒粗大，奥氏体及碳化物沿马氏体晶界分布，多相组织分布不均匀。
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；从图上可以看出，采用本发明方法，可获得非常均匀、细小、弥散分布的各相组织。因此本发明的热镀锌Q&P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，进而改善了材料组织，提高了材料性能。
Figure PCTCN2022084518-appb-000019
Figure PCTCN2022084518-appb-000020
Figure PCTCN2022084518-appb-000021
Figure PCTCN2022084518-appb-000022
表19
Figure PCTCN2022084518-appb-000023
表20
Figure PCTCN2022084518-appb-000024
Figure PCTCN2022084518-appb-000025
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热镀机组进行工艺改造，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。
以热镀锌Q&P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、热镀锌技术又有着巨大的开发价值，两者的结合必将会为热镀锌Q&P钢的开发和生产提供更大的空间。

Claims (14)

  1. 抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其化学成分质量百分比为：C：0.16～0.23％，Si：1.1～2.0％，Mn：1.6～3.0％，P≤0.015％，S≤0.005％，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，所述抗拉强度≥1180MPa的低碳低合金Q&P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40～85％；
    4)快速热处理
    冷轧后的钢板快速加热至770～845℃，所述快速加热采用一段式或两段式；采用一段式快速加热时加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以50～500℃/s的加热速率从550～625℃加热至770～845℃；之后进行均热，均热温度为770～845℃，均热时间为10～60s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃，随后以50～200℃/s的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s，随后以10～30℃/s的加热速率加热至300～470℃进行回火处理，回火时间10～60s；回火结束后以30～100℃/s的冷却速率冷却至室温；
    优选地，所述抗拉强度≥1180MPa的低碳低合金热镀锌Q&P钢通过以下方法制备得到：
    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770～845℃，所述快速加热采用一段式或两段式；采用一段式快速加热时，加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以30～500℃/s的加热速率从550～625℃加热至770～845℃；之后进行均热，均热温度：770～845℃，均热时间：10～60s；
    均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃，随后以50～200℃/s的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s，随后以10～30℃/s的加热速率加热至460～470℃进行配分处理，配分时间10～60s；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10～300℃/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≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于，所述抗拉强度≥1180MPa的低碳低合金低碳低合金Q&P钢和热镀锌Q&P钢中：C的含量范围选自0.17～0.23％、0.19～0.21％和0.18～0.21％；
    Si的含量范围选自1.1～1.7％、1.3～1.5％、1.4～2.0％和1.6～1.8％；
    Mn的含量范围选自1.6～2.2％、1.8～2.0％、2.4～3.0％和2.6～2.8％
    Cr的含量≤0.35％，如≤0.25％；
    Mo的含量≤0.25％；
    Nb的含量≤0.06％，如≤0.04％；
    Ti的含量≤0.065％，如≤0.04％，如0.006～0.016％；
    V的含量≤0.055％，如≤0.035％。
  3. 如权利要求1或2所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P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1-3中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于：
    所述的快速热处理全过程用时为71～186s，所述快速热处理和热镀锌全过程用时为43～186s；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770～845℃；优选地， 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770～845℃；和/或
    钢板的快速冷却速率为50～150℃/s。
  5. 如权利要求1-4中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于：
    所述低碳低合金Q&P钢的金相组织为马氏体75～90％、残余奥氏体10～25％、铁素体3～10％的多相组织，其基体组织分布均匀，具有片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相；和/或
    所述低碳低合金Q&P钢-50℃奥氏体转变率低于8％，-190℃奥氏体转变率低于30％；和/或
    所述低碳低合金Q&P钢的屈服强度≥660MPa，抗拉强度为≥1180MPa，延伸率≥18％，强塑积≥24GPa％；优选地，屈服强度为668～1112MPa，抗拉强度为1181～1350MPa，延伸率为18.9～24.2％，强塑积为24.1～28.97GPa％。
  6. 权利要求1-5中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于，所述低碳低合金Q&P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≤0.005％，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Q&P钢中，C含量为0.19～0.21％，和/或Si含量为1.3～1.5％，和/或Mn含量为1.8～2.0％；
    优选地，所述低碳低合金Q&P钢的金相组织为马氏体75～85％、残余奥氏体10～25％和铁素体3～10％的多相组织；
    优选地，所述低碳低合金Q&P钢的屈服强度为668～1002MPa，抗拉强度为1181～1296MPa，延伸率为18.9～24.2％，强塑积为24.1～28.6GPa％。
  7. 如权利要求1-5中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于，所述低碳低合金Q&P钢的化学成分质量百分比为：C：0.16～0.23％，Si：1.4～2.0％，Mn：2.4～3.0％，Ti：0.006～0.016％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、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Q&P钢的抗拉强度≥1280MPa；
    优选地，所述低碳低合金Q&P钢中，C含量为0.18～0.21％，和/或，Si含量为1.6～1.8％， 和/或，Mn含量为2.6～2.8％；
    优选地，所述低碳低合金Q&P钢的金相组织为马氏体80～90％、残余奥氏体10～20％、铁素体3～5％的多相组织；
    优选地，所述低碳低合金Q&P钢的屈服强度754～1112MPa，抗拉强度1281～1350MPa，延伸率19～22.2％，强塑积24.8～28.97GPa％。
  8. 如权利要求1-4中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于：
    所述热镀锌Q&P钢的金相组织为马氏体、铁素体和奥氏体三相组织，其基体组织分布均匀，具有片层状回火马氏体，晶粒粒径为1-3μm，马氏体强化相晶粒周围存在均匀分布的铁素体相；优选地，所述热镀锌Q&P钢的金相组织按体积分数占比为马氏体45～75％、铁素体15～30％、奥氏体10～25％的三相组织；和/或
    所述热镀锌Q&P钢的屈服强度≥720MPa，抗拉强度≥1180MPa，延伸率≥19％，强塑积≥23.0GPa％；优选地，所述热镀锌Q&P钢的屈服强度为721～956MPa，抗拉强度为1184～1352MPa，延伸率为19～22.5％，强塑积为23.6～28.9GPa％；和/或
    所述热镀锌Q&P钢金相组织-50℃奥氏体转化变率低于8％，-190℃奥氏体转变率低于30％。
  9. 如权利要求1-4和8中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于，所述低碳低合金热镀锌Q&P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≤0.005％，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，所述热镀锌Q&P钢中，C含量为0.19～0.21％，和/或，Si含量为1.3～1.5％，和/或，Mn含量为1.8～2.0％；
    优选地，所述热镀锌Q&P钢的金相组织中按体积分数占比为马氏体45～75％、铁素体15～30％和奥氏体10～25％的三相组织；
    优选地，所述热镀锌Q&P钢的屈服强度为721～805MPa，抗拉强度为1184～1297MPa，延伸率为19.1～22.4％，强塑积为23.6～28GPa％。
  10. 如权利要求1-4和8中任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢，其特征在于，所述低碳低合金热镀锌Q&P钢的化学成分质量百分比为：C：0.16～0.23％，Si：1.4～2.0％，Mn：2.4～3.0％，Ti 0.006～0.016％，P≤0.015％，S≤0.002％，Al：0.02～0.05％，还可含有Cr、Mo、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Q&P钢的抗拉强度≥1280MPa；
    优选地，所述热镀锌Q&P钢中，C含量为0.18～0.21％，和/或，Si含量为1.6～1.8％，和/或，Mn含量为2.6～2.8％；
    优选地，所述热镀锌Q&P钢的屈服强度为802～956MPa，抗拉强度为1280～1352MPa，延伸率19～22.5％，强塑积25.2～28.9GPa％。
  11. 如权利要求1～10任一项所述的抗拉强度≥1180MPa的低碳低合金Q&P钢或低碳低合金热镀锌Q&P钢的制造方法，其特征在于，所述低碳低合金Q&P钢的制造方法包括以下步骤：
    1)冶炼、铸造
    按所述的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40～85％，冷轧后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770～845℃，所述快速加热采用一段式或两段式；采用一段式快速加热时加热速率为50～500℃/s；采用两段式快速加热时，第一段以15～500℃/s的加热速率从室温加热至550～625℃，第二段以50～500℃/s的加热速率从550～625℃加热至770～845℃；
    b)均热
    在奥氏体和铁素体两相区目标温度770～845℃进行均热，均热时间为10～60s；
    c)冷却
    带钢或钢板均热结束后以5～15℃/s的冷却速率缓慢冷却至700～770℃，随后以50～200℃/s的冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s；
    d)回火
    保温结束后，将带钢或钢板以10～30℃/s的加热速率加热至300～470℃进行回火处理，回火时间10～60s；
    e)回火结束后带钢或钢板冷却至室温，冷却速率30～100℃/s；
    所述低碳低合金热镀锌Q&P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770～845℃奥氏体和铁素体两相区目标温度，所述快速加热采用一段式或两段式；
    采用一段式快速加热时，加热速率为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的加热速率从550～625℃加热至770～845℃；
    b)均热
    在奥氏体和铁素体两相区目标温度770～845℃进行均热，均热时间为10～60s；
    c)冷却
    带钢或钢板均热结束后以5～15℃/s冷却速率缓冷至720～770℃；随后以50～200℃/s冷却速率快速冷却至230～280℃，并在此温度区间保温2～10s(如2～8s)；
    d)配分
    保温结束后，将带钢或钢板以10～30℃/s的加热速率加热至460～470℃进行配分处理，配分时间10～60s；
    e)热镀锌
    配分结束后将带钢或钢板浸入锌锅进行热镀锌；
    f)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10～300℃/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的冷却速率快速冷却至室温，获得合金化热镀锌GA产品。
  12. 如权利要求11所述的方法，其特征在于，所述卷取温度为580～650℃。
  13. 如权利要求10或11所述的方法，其特征在于，所述冷轧压下率为60～80％。
  14. 如权利要求11～13中任一项所述的方法，其特征在于，
    所述低碳低合金Q&P钢的快速热处理全过程用时为71～186s，所述低碳低合金热镀锌 Q&P钢的快速热处理、热镀锌全过程用时为43～186s；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770～845℃；优选地，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770～845℃；和/或
    快速加热步骤中，所述快速加热最终温度为790～845℃；和/或
    冷却步骤中，所述带钢或钢板快速冷却速率为50～150℃/s；和/或
    均热过程中，带钢或钢板加热至所述奥氏体和铁素体两相区目标温度后，保持温度不变进行均热；和/或
    均热过程中，带钢或钢板在均热时间段内进行小幅度升温或小幅度降温，升温后温度不超过845℃，降温后温度不低于770℃；和/或
    所述均热时间为10～40s。
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CN117089761B (zh) * 2023-05-30 2024-06-04 宝山钢铁股份有限公司 一种归一化成分的变强度双相钢板及其柔性制造方法
CN116855829A (zh) * 2023-07-07 2023-10-10 天津市产品质量监督检测技术研究院检测技术研究中心 一种低碳纳米贝氏体钢及其制备方法
CN116855829B (zh) * 2023-07-07 2024-02-27 天津市产品质量监督检测技术研究院检测技术研究中心 一种低碳纳米贝氏体钢及其制备方法
CN116949366B (zh) * 2023-08-01 2024-03-19 北京理工大学 无碳化物的回火马氏体钢及其制备方法
CN116949366A (zh) * 2023-08-01 2023-10-27 北京理工大学 无碳化物的回火马氏体钢及其制备方法
US12215404B1 (en) 2023-08-01 2025-02-04 Beijing Institute Of Technology Carbide-free tempered martensitic steels and preparation methods thereof
CN117144259A (zh) * 2023-08-07 2023-12-01 武汉钢铁有限公司 一种适于-40℃环境使用的hb450级耐磨钢及生产方法
CN118460914A (zh) * 2024-04-19 2024-08-09 首钢集团有限公司 一种780MPa级合金化热镀锌双相钢及其制备方法

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