Low-carbon, low-alloy, highly formable duplex steel with tensile strength ≥ 590 MPa, hot-dip galvanized duplex steel, and method for manufacturing the same.

Abstract

Provided is a low carbon low alloy highly formable duplex steel or highly formable hot-dip galvanized duplex steel having a tensile strength of ≥ 590MPa and a rapid heat treatment manufacturing method thereof, the chemical components of the steel are, in mass percentage, C: 0.04-0.12%, Si: 0.1-0.5%, Mn: 1.0-2.0%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02-0.06%, and may further contain one or two of Cr, Mo, Ti, Nb, and V, where Cr + Mo + Ti + Nb + V ≤ 0.5%, with the balance being Fe and unavoidable impurities. The manufacturing method includes the steps of smelting, casting, hot rolling, cold rolling, and rapid heat treatment. By controlling the rapid heating, short-time heating and rapid cooling processes in the rapid heat treatment process, the present invention changes the recovery of the deformed structure, recrystallization and austenite transformation processes, increases the nucleation rate, shortens the grain growth time, refines the grains, increases the strength and n value of the material, and expands the range of the material performance range.

Description

[Technical Field]

[0001] The present invention belongs to the field of rapid heat treatment technology for materials, and more particularly to low-carbon, low-alloy, highly formable duplex steel and highly formable hot-dip galvanized duplex steel having a tensile strength of ≥ 590 MPa, and a method for rapidly heat-treating and manufacturing them. [Background technology]

[0002] As public awareness of energy conservation and the safety of materials and services gradually increases, many automakers are selecting high-strength steel sheets as automotive materials; in the automotive industry, high-strength steel sheets are used to reduce the thickness of the steel sheets while improving the dent resistance, durability, large deformation impact toughness, and crash safety of vehicles. Therefore, automotive steel sheets are inevitably evolving in the direction of high strength, high toughness, and easy formability.

[0003] Among high-strength steels for automotive applications, duplex steel is the most widely used and has the best application prospects. Low-carbon, low-alloy duplex steel is characterized by a low yield ratio, a high initial work hardening rate, and a good synergy of strength and plasticity, making it widely used as a press steel for automotive structures due to its high strength and good formability.

[0004] Conventional duplex steel is obtained by cold-rolling low-carbon steel or low-alloy high-strength steel, soaking and annealing it in the critical region, followed by rapid cooling or hot rolling, controlled rolling, and controlled cooling. Its microstructure mainly consists of ferrite and martensite. Duplex steel utilizes the principle of "composite materials" to maximize the advantages of each phase (ferrite and martensite) in the steel, while mitigating or eliminating the disadvantages of one phase through the presence of the other phase.

[0005] The mechanical properties of duplex steel depend primarily on the following three aspects: 1. Size of the crystal grains and distribution of alloying elements in the matrix phase; II. Size, shape, distribution, and volume fraction of the second phase; 3. The characteristic of both the matrix and the second phase binding together.

[0006] Therefore, how to acquire duplex steel products that are low-cost, high-performance, and possess good strength-plasticity synergies has become a goal pursued by each steel company, attracting widespread interest from steel companies and automotive users.

[0007] Cold-rolled duplex steel is obtained by a rapid cooling process after soaking to a critical temperature, and this process mainly consists of three steps: Step 1: Heat the steel strip to the critical temperature of the ferrite and austenite phases, and maintain the temperature evenly; Step 2: Cool the sample at a cooling rate higher than the critical cooling rate. s ~M f The mixture is cooled to a certain temperature to obtain a two-phase structure consisting of a certain amount of martensite and ferrite; Step 3: Keep the steel strip warm, or M s The material is heated to the following temperature, kept warm, and then tempered to obtain a good microstructural synergy between the hard martensite phase and the soft ferrite phase, ultimately resulting in a two-phase structure of martensite and ferrite.

[0008] Currently, 590 MPa class cold-rolled duplex steel produced by conventional continuous annealing methods has a slow heating rate, and both the heating time and soaking time are relatively long, so the entire continuous annealing cycle takes 5-8 minutes; the recrystallization and transformation processes during the heating process occur sequentially and generally do not overlap, so the ferrite recrystallized grains and austenite grains each undergo nucleation and sufficient growth, and finally the size of the resulting ferrite and martensite duplex grain structure is relatively large, usually around 10-20 μm.

[0009] The primary control methods for duplex steel in conventional technology involve adding alloying elements and adjusting the soaking temperature in the annealing process, as well as the temperature and time of the quenching and tempering processes, thereby altering the phase structure ratio and distribution of the duplex steel and obtaining relatively optimized product performance.

[0010] Chinese Patent Application CN109 1Document 61805A discloses "590MPa-class lightweight cold-rolled duplex steel for automobiles and its production method." The chemical composition of the steel of this invention is, by weight percentage, C: 0.06~0.08%, Si: 0.05~0.10%, Mn: 1.7~1.8%, P≦0.01%, S≦0.005%, Al: 0.02~0.05%, and Cr: 0.20~0.30%. It was obtained by a conventional continuous annealing method. The mechanical properties of the duplex steel are a yield strength of 374-406 MPa, a tensile strength of 630-653 MPa, and an elongation of 24-26%. The relatively high alloy content in the steel of this invention increases costs, makes the manufacturing process difficult, and reduces usability such as material welding performance.

[0011] Chinese patent application CN106011643B discloses "Cold-rolled duplex steel with a tensile strength of 590 MPa and method for manufacturing it." The chemical composition of the steel of this invention is, by weight percentage, C: 0.06~0.1%, Si: 0.26~0.6%, Mn: 1.2~1.6%, Cr: 0.10~0.45%, Al: 0.02~0.06%, P≦0.02%, S≦0.015%, N≦0.006%, with the remainder being Fe and other unavoidable impurities. The method of this invention employs a conventional continuous annealing method, and a small amount of Cr is added instead of Mn to ensure hardenability. The mechanical properties of the duplex steel of this invention are a yield strength of 290-330 MPa, a tensile strength of 600-650 MPa, and an elongation of 23-27%. In this invention, the element Cr, a precious metal, is added to reduce the Mn content, which increases both the manufacturing cost and the difficulty of manufacturing. Furthermore, because the alloy content is relatively high, the annealing temperature is also high.

[0012] Chinese Patent Application CN105543674B discloses a "method for manufacturing a cold-rolled ultra-high-strength dual-phase steel with high local forming performance". The chemical composition of the high-strength dual-phase steel of this invention is, by weight percentage, C: 0.08 - 0.12%, Si: 0.1 - 0.5%, Mn: 1.5 - 2.5%, Al: 0.015 - 0.05%, and the balance is Fe and other inevitable impurities. Selectively blend raw materials according to this chemical composition and melt them into billets; heat the billets at 1150 - 1250°C for 1.5 - 2 hours, then perform hot rolling. The starting temperature of hot rolling is 1080 - 1150°C, and the ending temperature of rolling is 880 - 930°C; after rolling, cool to 450 - 620°C at a cooling rate of 50 - 200°C / s, perform coiling, and obtain a hot-rolled steel sheet with bainite as the main tissue type; perform cold rolling on the hot-rolled steel sheet, then heat it to 740 - 820°C at a rate of 50 - 300°C / s, perform annealing, keep the temperature for 30s - 3min, cool to 620 - 680°C at a cooling rate of 2 - 6°C / s, then cool to 250 - 350°C at a cooling rate of 30 - 100°C / s, and perform overaging treatment for 3 - 5min to obtain an ultra-high-strength dual-phase steel with a ferrite + martensite dual-phase structure. The yield strength of this ultra-high-strength dual-phase steel is 650 - 680 MPa, the tensile strength is 1023 - 1100 MPa, the elongation is 12.3 - 13%, and it does not crack even when bent 180 o degrees along the rolling direction.

[0013] The most main feature of this patent is to combine the control of the cooling condition after hot rolling and the rapid heating in the continuous annealing process, that is, by controlling the cooling process after hot rolling, strip structure is removed to achieve the homogenization of the structure; during the subsequent continuous annealing process, rapid heating is adopted to ensure the uniformity of the structure and achieve the refinement of the structure. The premise for this patent technology to adopt rapid heating annealing is to obtain a hot-rolled raw material with bainite as the main structure after hot rolling, and its purpose is mainly to ensure the uniformity of the structure and avoid local deformation and non-uniformity caused by strip structure.

[0014] The disadvantages of this patent are mainly as follows: First, it is necessary to obtain a hot-rolled raw material with a bainite structure; this hot-rolled raw material has high strength and large deformation resistance, bringing great difficulties to subsequent pickling and cold rolling production. Second, the understanding of rapid heating is limited to shortening the heating time and refining the crystal grains, and its heating rate is not classified according to the changes in the material microstructure at different temperature stages. Instead, it is heated at a speed of 50 - 300 °C / s for all, resulting in an increase in the production cost of rapid heating. Third, the soaking time is 30 s - 3 min; an increase in the soaking time will inevitably weaken the crystal grain refinement effect caused by rapid heating to some extent, which is disadvantageous for improving the strength and toughness of the material. Fourth, this method requires an overaging treatment for 3 - 5 minutes, but this is actually too long for the aging time of rapid heat-treated DP steel and is not necessary. Also, an increase in both the soaking time and the overaging time is disadvantageous for energy conservation, reduction of mechanical equipment investment and machine floor area, and is also disadvantageous for the high-speed and stable operation of the strip steel in the furnace. It is obvious that this is not a rapid heat treatment process in a strict sense.

[0015] Chinese Patent Application 201711385126.5 discloses a "780 MPa grade low-carbon low-alloy TRIP steel", and its chemical composition is, by mass percentage, C: 0.16 - 0.22%, Si: 1.2 - 1.6%, Mn: 1.6 - 2.2%, and the balance is Fe and inevitable impurities, which is obtained by the following rapid heat treatment process: The strip steel is rapidly heated from room temperature to the austenite and ferrite two-phase region of 790 - 830 °C, and the heating rate is 40 - 300 °C / s; the residence time in the two-phase region heating target temperature range is 60 - 100 s; the strip steel is rapidly cooled from the two-phase region temperature to 410 - 430 °C, and the cooling rate is 40 - 100 °C / s, and it stays in this temperature range for 200 - 300 s; the strip steel is rapidly cooled from 410 - 430 °C to room temperature. Its characteristics are that the metal microstructure of the said TRIP steel is a three-phase structure of bainite, ferrite, and austenite; the average crystal grain size of the said TRIP steel is clearly refined; the tensile strength is 950 - 1050 MPa; the elongation rate is 21 - 24%; the strength-elongation product can reach a maximum of 24 GPa%.

[0016] The shortcomings of this patent are mainly as follows: Firstly, while the patent discloses a 780 MPa class low-carbon low-alloy TRIP steel product and its process technology, the tensile strength of the TRIP steel product is 950-1050 MPa. This strength is too high for a 780 MPa class product, resulting in unfavorable user performance, but it is also too low for a 980 MPa class product, failing to adequately meet user strength requirements. Secondly, the patent employs a single-stage rapid heating method, using the same rapid heating rate throughout the entire heating temperature range. It does not differentiate between different temperature stages based on changes in the material's microstructure; instead, it rapidly heats everything at a rate of 40-300°C / s, which inevitably leads to increased production costs for the rapid heating process. Thirdly, the soaking time in this patent is set to 60-100 seconds, which is close to the soaking time of conventional continuous annealing; increasing the soaking time inevitably weakens the grain refinement effect due to rapid heating, which is very detrimental to improving material strength and toughness; Fourth, the patent requires a bainite isothermal treatment time of 200-300 seconds, which is actually too long for rapidly heat-treated products, preventing them from fulfilling their intended role and making it unnecessary. Furthermore, increasing both the soaking time and the isothermal treatment time is detrimental to energy saving, reduction of machinery investment and machine footprint, and also detrimental to the high-speed stable operation of the steel strip in the furnace, and is clearly not a rapid heat treatment process in the strict sense.

[0017] Chinese patent application CN108774681A discloses a "rapid heat treatment method for high-strength steel," which employs a ceramic sheet electric heating device to achieve a heating rate reaching a maximum of 400°C / s. After heating to 1000-1200°C, the steel is cooled by blowing air with a fan, and then cooled to room temperature at a maximum cooling rate close to 3000°C / s. In this invention, the processing speed of the heat treatment apparatus using electric heating of the ceramic sheet is 50 cm / min. The steel according to this invention is characterized by a high carbon content of 0.16-0.55% and also contains alloying elements such as Si, Mn, Cr, and Mo; this method is mainly suitable for steel wire, wire coils, or steel strips of 5 mm or less. The patent describes a rapid heat treatment method using electric heating of ceramic sheets; the main objective of the invention is to solve the problems of energy waste and environmental pollution caused by the low heat treatment efficiency of products such as high-strength steel wire and wire coils; it does not mention the effects and actions of rapid heating on the microstructure performance of the material; the invention employs a blower-assisted cooling method regardless of the steel grade number, composition, and microstructure characteristics; however, the fact that the maximum cooling rate is close to 3000°C / s should refer to the instantaneous cooling rate in the high-temperature segment, and as is well known, the average cooling rate of blow-assisted cooling cannot reach 3000°C / s; at the same time, producing wide, thin steel strips using excessively high cooling rates in the high-temperature segment leads to excessive internal stress and problems such as poor sheet shape in the steel plate, making it unsuitable for large-scale industrialized continuous heat treatment production of wide, thin steel plates.

[0018] Chinese patent application CN106811698B discloses "High-strength steel sheet with precise control of structure and method for manufacturing the same," wherein the chemical composition of the high-strength duplex steel is, by weight percentage, C: 0.08~0.40%, Si: 0.35~3.5%, Mn: 1.5~7.0%, P: ≤0.02%, S: ≤0.02%, Al: 0.02~3.0%, and further includes at least one of Cr: 0.50~1.5%, Mo: 0.25~0.60%, Ni: 0.5~2.5%, Cu: 0.20~0.50%, B: 0.001~0.005%, V: 0.10~0.5%, Ti: 0.02~0.20%, Nb: 0.02~0.20%, with the remainder being Fe and other unavoidable impurities. Its mechanical properties: Tensile strength R mexceeds 1000 MPa, and the elongation rate A 50mm exceeds 28%. In the present invention, the contents of components C, Si, and Mn are all relatively high. By performing non-isothermal annealing on a conventional continuous annealing production line and using a method of removing the soaking section, steel strips of different components are recrystallized annealed. The specific range of annealing parameters is that after rapidly heating to 800 - 930 °C at 20 °C / s or more, immediately cooling at a cooling rate of 40 °C / s or more to the M s -M f point, and then maintaining the temperature at the M f ~M f point + 100 °C for 30 s to 30 min, and finally cooling to room temperature.

[0019] The main feature of the present invention is to control the morphology and structure of the martensite strengthening phase to obtain a fine martensite structure with fine needle-like and short rod-like shapes. By reheating, C atoms are diffused into the retained austenite, and finally relatively stable retained austenite is obtained, endowing it with a certain deformation ability, thereby enhancing the plasticity and toughness of the high-strength steel.

[0020] The so-called rapid heating in this invention actually has a low heating rate, the heating rate is 20 - 60 °C / s, belonging to a medium heating rate, and the cooling rate is 40 - 100 °C / s. The considerations for rapid heating, rapid cooling, and omission of the soaking section are to shorten the residence time of the high-strength steel at the high-temperature section, ensure that the grain size in the austenitization process of the steel is fine and the structure and chemical composition are not completely homogenized, and guarantee that a large amount of large-sized lath-like martensite is not generated after cooling, and to obtain a certain amount of film-like retained austenite structure. However, this inevitably makes it difficult to control the heating temperature and increases the fluctuations in the microstructure and properties.

[0021] This method, while still based on conventional heating and cooling techniques using continuous annealing units, omits the soaking segment (reducing the soaking time to zero), increases the alloy content, and after quenching and tempering, ultimately obtains a high-strength steel product with a synergy of strength and toughness. However, this invention has not been specifically subdivided and researched for each strength level of steel grade number. Furthermore, the heating rate is moderate, not rapid heating, and there is no soaking time, so it does not embody a truly rapid heat treatment method or a complete annealing cycle, and therefore has no future potential for commercial application.

[0022] Chinese patent application CN107794357B and US patent application US2019 / 0153558A1 describe a method for producing ultra-high-strength martensitic cold-rolled steel sheets by an ultra-rapid heating process. The chemical composition of the high-strength duplex steel is, 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%, with the remainder being Fe and other unavoidable impurities. Mechanical properties of the duplex steel: Yield strength Rp 0.2 It exceeds 1100 MPa, tensile strength R m The pressure is 1800-2300 MPa, the elongation is a maximum of 12.3%, and the uniform elongation is 5.5-6%. The present invention provides an ultra-rapid heating production process for ultra-high-strength martensitic cold-rolled steel sheets, characterized by first heating the cold-rolled steel sheet to 300-500°C at a rate of 1-10°C / s, and then reheating it to the single-phase austenite region of 850-950°C at a heating rate of 100-500°C / s; after that, the steel sheet is kept warm for no more than 5 seconds and then immediately cooled to room temperature with water to obtain an ultra-high-strength cold-rolled steel sheet.

[0023] The drawbacks of the process described in the patent include: Firstly, the annealing temperature of the steel in this invention falls within the ultra-high temperature range of the austenite single-phase region, and moreover, it contains many alloying elements, and both the yield strength and tensile strength exceed 1000 MPa. This has caused significant difficulties in the heat treatment process, the pre-heat treatment manufacturing process, and subsequent use by users.

[0024] Secondly, the ultra-rapid heating annealing method of the present invention employs a holding time of no more than 5 seconds, which not only results in poor controllability of the heating temperature but also leads to uneven distribution of alloying elements in the final product, as well as uneven and unstable product microstructure performance; Thirdly, the final rapid cooling involved water quenching to room temperature without the necessary tempering treatment. As a result, the resulting final product microstructure and the alloy element distribution profile within the final microstructure did not provide optimal toughness, leaving the final product with excessive strength and insufficient plasticity and toughness. Fourthly, the method of the present invention involves a cooling rate that is too high due to water quenching, which causes problems such as deformation of the steel plate and surface oxidation. Therefore, the patented technology has no practical value or only low practical value.

[0025] Currently, conventional continuous annealing furnace production lines are limited by their equipment capacity. Research on cold-rolled duplex steel products and the annealing process involves slow heating of the strip steel based on the heating rate of existing industrial equipment (5-20°C / s), followed by sequential recovery, recrystallization, and austenitization transformation. This results in relatively long heating and soaking times and high energy consumption. Furthermore, conventional continuous annealing production lines have problems such as the long time the strip steel stays in the high-temperature furnace segment and the large number of rollers it passes through. Conventional continuous annealing units, based on product specifications and production capacity requirements, generally have a soaking time of 1-3 minutes and a unit speed of about 180 meters / minute. However, the number of rollers in the high-temperature furnace segment is generally 20-40, increasing the difficulty of controlling the quality of the strip steel surface. [Overview of the project] [Problems that the invention aims to solve]

[0026] The object of the present invention is to provide a low-carbon, low-alloy, highly formable duplex steel with a tensile strength of ≥590 MPa, a highly formable hot-dip galvanized duplex steel, and a rapid heat treatment method for manufacturing the same. Rapid heat treatment alters the recovery of the deformation structure, recrystallization, and austenite transformation processes, thereby improving the nucleation rate (including the recrystallization nucleation rate and the austenite transformation nucleation rate), shortening the grain growth time, and refining the grains. The resulting duplex steel has a yield strength of ≥340 MPa, a tensile strength of ≥620 MPa, an elongation of ≥19%, a strength-to-elongation product of ≥15.5 GPa%, and a strain hardening index of n. 90 The value is greater than 0.20; it improves material strength and provides good plasticity and toughness; the rapid heat treatment process improves production efficiency, reduces production costs and energy consumption, significantly reduces the number of furnace rollers, and improves the surface quality of the steel sheet. [Means for solving the problem]

[0027] To achieve the above objectives, the technical solution of the present invention is: A low-carbon, low-alloy, highly formable duplex steel or a low-carbon, low-alloy, highly formable hot-dip galvanized duplex steel with a tensile strength of ≥ 590 MPa, the chemical composition of which is, by mass percentage, C: 0.04~0.12%, Si: 0.1~0.5%, Mn: 1.0~2.0%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02~0.06%, and may further contain one or two from Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤ 0.5%, and the remainder being Fe and unavoidable impurities.

[0028] In some embodiments, the carbon content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 0.045-0.105%, 0.04-0.10%, or 0.05-1.2%. In some embodiments, the carbon content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 0.06-0.08%. In some embodiments, the carbon content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 0.065-0.085%. In some embodiments, the carbon content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 0.07-0.1%.

[0029] In some embodiments, the Si content of the highly formable duplex steel or highly formable hot-dip galvanized duplex steel is 0.1 to 0.4% or 0.1 to 0.3%. In some embodiments, the Si content of the highly formable duplex steel or highly formable hot-dip galvanized duplex steel is 0.15 to 0.25%.

[0030] In some embodiments, the Mn content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 1.0-1.6%, 1.0-1.5%, or 1.2-2.0%. In some embodiments, the Mn content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 1.2-1.4%. In some embodiments, the Mn content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 1.2-1.35%. In some embodiments, the Mn content of the high formability duplex steel or high formability hot-dip galvanized duplex steel is 1.5-1.8%.

[0031] In some embodiments, the highly formable duplex steel or highly formable hot-dip galvanized duplex steel has a Cr+Mo+Ti+Nb+V ≤ 0.3%.

[0032] In some embodiments, the chemical composition of the low-carbon, low-alloy, highly formable duplex steel having a tensile strength of ≥ 590 MPa is, by mass percentage, C: 0.04-0.105%, Si: 0.1-0.4%, Mn: 1.0-1.6%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02-0.06%, and may further contain one or two of Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤ 0.5%, and the remainder being Fe and unavoidable impurities.

[0033] In some embodiments, the low-carbon, low-alloy, highly formable duplex steel having a tensile strength of ≥ 590 MPa is obtained by the following process: 1) Smelting and casting The slab is smelted and cast according to the above chemical composition; 2) Hot rolling, coiling Hot rolling end temperature ≥ A r3 Winding temperature 550~680℃; 3) Cold rolling The reduction ratio in cold rolling is 40-85%; 4) Rapid heat treatment Cold-rolled steel sheets are rapidly heated to 750-845°C, using either a one-stage or two-stage heating method; when using a one-stage rapid heating method, the heating rate is 50-500°C / s; when using a two-stage rapid heating method, the first stage involves heating from room temperature to 550-650°C at a heating rate of 15-500°C / s, and the second stage involves heating from 550-650°C to 750-845°C at a heating rate of 30-500°C / s (e.g., 50-500°C / s); after which, soaking is performed, with a soaking temperature of 750-845°C and a soaking time of 10-60 seconds; After soaking is complete, the material is slowly cooled to 670-770°C at a cooling rate of 5-15°C / s, and then rapidly cooled from 670-770°C to room temperature at a cooling rate of 50-200°C / s. Alternatively, the material may be rapidly cooled from 670-770°C to 230-280°C at a cooling rate of 50-200°C / s, overaging treatment may be performed in this temperature range, with an overaging treatment time of 200s or less, and finally, the material may be cooled to room temperature at a cooling rate of 30-50°C / s.

[0034] Preferably, the total processing time for the rapid heat treatment described in step 4) is 30 to 300 seconds. In some embodiments, the total processing time for the rapid heat treatment described in step 4) is 41 to 300 seconds. In some embodiments, the total processing time for the rapid heat treatment described in step 4) is 41 to 296 seconds.

[0035] Preferably, in step 2), the winding temperature is 580 to 650°C. Preferably, in step 3), the reduction ratio of the cold rolling is 60-80%.

[0036] Preferably, in step 4), when a single-stage heating method is used for the rapid heating, the heating rate is 50 to 300°C / s.

[0037] Preferably, in step 4), a two-stage heating method is employed for the rapid heating: in the first stage, the temperature is heated from room temperature to 550 to 650°C at a heating rate of 15 to 500°C / s, and in the second stage, the temperature is heated from 550 to 650°C to 750 to 845°C at a heating rate of 50 to 300°C / s.

[0038] Preferably, in step 4), a two-stage heating method is employed for the rapid heating: in the first stage, the temperature is heated from room temperature to 550 to 650°C at a heating rate of 50 to 300°C / s, and in the second stage, the temperature is heated from 550 to 650°C to 750 to 845°C at a heating rate of 80 to 300°C / s.

[0039] Preferably, in step 4), the soaking time is 10 to 40 seconds. Preferably, in step 4), the rapid cooling rate of the steel plate is 50 to 150°C / s.

[0040] Preferably, the overaging treatment time is 20 to 200 seconds. The microstructure of the duplex steel or hot-dip galvanized duplex steel described in the present invention is a ferrite-martensite duplex structure with an average grain size of 2 to 10 μm. In some embodiments, the microstructure of the duplex steel or hot-dip galvanized duplex steel described in the present invention is a ferrite-martensite duplex structure with an average grain size of 2 to 8 μm or 4 to 10 μm.

[0041] In some embodiments, the yield strength of the duplex steel or hot-dip galvanized duplex steel described in the present invention is 340-560 MPa, the tensile strength is 620-880 MPa, the elongation is 19-30.5%, the strength-to-elongation product is 15.5-20.5 GPa%, and the strain hardening index is n 90 The value is greater than 0.20.

[0042] In some embodiments, the yield strength of the duplex steel described in the present invention is 340-410 MPa, the tensile strength is 640-710 MPa, the elongation is 22-30.5%, the strength-to-elongation product is 15.5-20.5 GPa%, and the strain-hardening index is n 90The value is greater than 0.20.

[0043] In some embodiments, the chemical composition of the low-carbon, low-alloy, highly formable duplex steel with a tensile strength of ≥590 MPa described in the present invention is, by mass percentage, C: 0.04~0.10%, Si: 0.1~0.3%, Mn: 1.0~1.6%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02~0.06%, and may further contain one or two from Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤ 0.5%, and the remainder being Fe and unavoidable impurities. Preferably, the C content of the duplex steel is 0.06~0.08%; preferably, the Si content of the duplex steel is 0.15~0.25%; preferably, the Mn content of the duplex steel is 1.2%~1.4%. Preferably, the microstructure of the duplex steel is a ferrite-martensite duplex structure with an average grain size of 4 to 10 μm. Preferably, the yield strength of the duplex steel is 350 to 410 MPa, the tensile strength is 620 to 710 MPa, the elongation is 24.0 to 30.5%, the strength-to-elongation product is 17 to 20.5 GPa%, and the strain hardening index is n 90 The value is greater than 0.20; preferably, the yield strength of the duplex steel is 350-405 MPa, the tensile strength is 624-706 MPa, the elongation is 24.4-30.4%, the strength-to-elongation product is 17-20.1 GPa%, and the strain hardening index is n 90 The value is greater than 0.20. Preferably, the duplex steel is obtained by manufacturing it using the method of any one of the embodiments specified herein. Preferably, in the manufacturing process of the duplex steel, the total process time of the rapid heat treatment described in step 4) is 41 to 300 s.

[0044] In some embodiments, the chemical composition of the low-carbon, low-alloy, highly formable duplex steel with a tensile strength of ≥590 MPa described in the present invention is, by mass percentage, C: 0.05~0.12%, Si: 0.01~0.5%, Mn: 1.2~2.0%, P ≤0.015%, S ≤0.003%, Al: 0.02~0.055%, and may further contain one or two from Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤0.5%, and the remainder being Fe and unavoidable impurities. Preferably, the tensile strength of the duplex steel is ≥780 MPa. Preferably, the C content of the duplex steel is 0.07~0.1%; preferably, the Si content of the duplex steel is 0.1~0.4%; preferably, the Mn content of the duplex steel is 1.5~1.8%. Preferably, the microstructure of the duplex steel is a ferrite-martensite duplex structure with an average grain size of 2 to 8 μm. Preferably, the yield strength of the duplex steel is 400 to 540 MPa, the tensile strength is 780 to 880 MPa, the elongation is 19 to 24.5%, the strength-to-elongation product is 16.0 to 19.5 GPa%, and the strain hardening index is n 90 The value is greater than 0.20. Preferably, the yield strength of the duplex steel is 400-533 MPa, the tensile strength is 781-878 MPa, the elongation is 19.5-24.1%, the strength-to-elongation product is 16.3-19.3 GPa%, and the strain hardening index is n 90 The value is greater than 0.20. Preferably, the duplex steel is obtained by manufacturing it by the method of any one of the embodiments specified herein. Preferably, in the manufacturing process of the duplex steel, the total process time of the rapid heat treatment described in step 4) is 41 to 296 s.

[0045] In some embodiments, the chemical composition of the low-carbon, low-alloy, highly formable hot-dip galvanized duplex steel having a tensile strength of ≥ 590 MPa is, by mass percentage, C: 0.045-0.12%, Si: 0.1-0.5%, Mn: 1.0-2.0%, P ≤ 0.02%, S ≤ 0.006%, Al: 0.02-0.055%, and may further contain one or two of Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤ 0.5%, and the remainder being Fe and unavoidable impurities.

[0046] In some embodiments, the duplex steel is a hot-dip galvanized duplex steel described in any one of the embodiments specified herein, wherein after the soaking step is completed and the steel is slowly cooled to 670-770°C, it is rapidly cooled to 460-470°C at a cooling rate of 50-200°C / s, and then immersed in a zinc pot for hot-dip galvanizing to obtain the hot-dip galvanized duplex steel. In some embodiments, after hot-dip galvanizing, the steel is rapidly cooled to room temperature at a cooling rate of 30-150°C / s to obtain a hot-dip pure galvanized GI product. In some embodiments, after hot-dip galvanizing, the steel is heated to 480-550°C at a heating rate of 30-200°C / s for an alloying treatment, with an alloying treatment time of 10-20 s; after the alloying treatment, the steel is rapidly cooled to room temperature at a cooling rate of 30-250°C / s to obtain an alloyed hot-dip galvanized GA product.

[0047] Preferably, in some embodiments, the chemical composition of the hot-dip galvanized duplex steel is, by mass percentage, C: 0.045-0.105%, Si: 0.1-0.4%, Mn: 1.0-1.5%, P ≤ 0.02%, S ≤ 0.006%, Al: 0.02-0.055%, and may further contain one or two from Cr, Mo, Ti, Nb, and V, with Cr+Mo+Ti+Nb+V ≤ 0.3%, and the remainder being Fe and unavoidable impurities. Preferably, in the hot-dip galvanized duplex steel, the C content is 0.065-0.085%; preferably, in the hot-dip galvanized duplex steel, the Si content is 0.15-0.25%; preferably, in the hot-dip galvanized duplex steel, the Mn content is 1.2%-1.35%. Preferably, in the hot-dip galvanized duplex steel, Cr+Mo+Ti+Nb+V ≤ 0.2%. Preferably, the yield strength of the hot-dip galvanized duplex steel is 340-410 MPa, the tensile strength is 620-710 MPa, the elongation is 22-30.5%, the strength-to-elongation product is 15.5-20.0 GPa%, and the strain hardening index is n 90 The value is greater than 0.21; more preferably, the yield strength of the hot-dip galvanized duplex steel is 347-405 MPa, the tensile strength is 624-709 MPa, the elongation is 22.2-30.3%, the strength-to-elongation product is 15.7-19.6 GPa%, and the strain hardening index is n 90 The value is greater than 0.21.

[0048] In some embodiments, the chemical composition of the hot-dip galvanized duplex steel is, by mass percentage, C: 0.05-0.12%, Si: 0.01-0.5%, Mn: 1.2-2.0%, P ≤ 0.015%, S ≤ 0.003%, Al: 0.02-0.055%, and may further contain one or two from Cr, Mo, Ti, Nb, and V, with Cr + Mo + Ti + Nb + V ≤ 0.5%, and the remainder being Fe and unavoidable impurities. Preferably, the yield strength of the hot-dip galvanized duplex steel is ≥ 780 MPa. Preferably, the C content of the hot-dip galvanized duplex steel is 0.07-0.1%; preferably, the Si content of the hot-dip galvanized duplex steel is 0.1-0.4%; preferably, the Mn content of the hot-dip galvanized duplex steel is 1.5-1.8%. Preferably, the hot-dip galvanized duplex steel has a yield strength of 400-560 MPa, a tensile strength of 790-870 MPa, an elongation of 19.0-25.0%, a strength-to-elongation product of 16.0-20.0 GPa%, and a strain hardening index of n. 90 The value is greater than 0.20; more preferably, the yield strength of the hot-dip galvanized duplex steel is 400-552 MPa, the tensile strength is 798-862 MPa, the elongation is 19.5-24.6%, the strength-to-elongation product is 16.3-19.9 GPa%, and the strain hardening index is n 90 The value is greater than 0.20.

[0049] In some embodiments, the highly formable hot-dip galvanized duplex steel described in the present invention is obtained by the following process: a) Smelting, casting The slab is smelted and cast according to the above chemical composition; b) Hot rolling, coiling Hot rolling end temperature ≥ A r3 Winding temperature 550~680℃; c) Cold rolling The reduction ratio in cold rolling is 40-85%; d) Rapid heat treatment, hot-dip galvanizing Cold-rolled steel sheets are rapidly heated to 750-845°C, and either a one-stage or two-stage heating method is employed for this rapid heating process. When employing a single-stage rapid heating system, the heating rate is 50-500°C / s; When employing a two-stage rapid heating system, the first stage involves heating from room temperature to 550-650°C at a heating rate of 15-500°C / s, and the second stage involves heating from 550-650°C to 750-845°C at a heating rate of 30-500°C / s (e.g., 50-500°C / s); after which, soaking is performed, with a soaking temperature of 750-845°C and a soaking time of 10-60 seconds. After soaking, the material is slowly cooled to 670-770°C at a cooling rate of 5-15°C / s, then rapidly cooled to 460-470°C at a cooling rate of 50-200°C / s (e.g., 50-150°C / s), immersed in a zinc pot, and hot-dip galvanizing is performed. After hot-dip galvanizing, the product is rapidly cooled to room temperature at a cooling rate of 30-150°C / s to obtain a hot-dip pure zinc-plated GI product. Alternatively, after hot-dip galvanizing, the material is heated to 480-550°C at a heating rate of 30-200°C / s to perform an alloying treatment, with an alloying treatment time of 10-20 seconds; 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.

[0050] Preferably, the total processing time for the rapid heat treatment and hot-dip galvanizing described in step d) is 30 to 142 seconds.

[0051] Preferably, in step b), the winding temperature is 580 to 650°C. Preferably, in step c), the reduction ratio of the cold rolling is 60-80%.

[0052] Preferably, in step d), when a single-stage heating method is used for the rapid heating, the heating rate is 50 to 300°C / s.

[0053] Preferably, in step d), a two-stage heating method is employed for the rapid heating: in the first stage, the temperature is heated from room temperature to 550 to 650°C at a heating rate of 15 to 500°C / s, and in the second stage, the temperature is heated from 550 to 650°C to 750 to 845°C at a heating rate of 50 to 300°C / s.

[0054] Preferably, in step d), a two-stage heating method is employed for the rapid heating: in the first stage, the temperature is heated from room temperature to 550 to 650°C at a heating rate of 30 to 500°C / s, and in the second stage, the temperature is heated from 550 to 650°C to 750 to 845°C at a heating rate of 80 to 300°C / s.

[0055] The metallic structure of the hot-dip galvanized duplex steel described in the present invention is a uniformly distributed ferrite and martensite duplex structure, and the average grain size is 1 to 3 μm.

[0056] In the design of the steel composition and process according to the present invention: C: Carbon is the most common strengthening element in steel. While carbon increases the strength of steel and decreases its plasticity, the carbon content should not be too high, as low yield strength, high uniform elongation, and total elongation are necessary for forming steel. The carbon content has a significant impact on the mechanical properties of steel. As the carbon content increases, the number of pearlites increases, greatly improving the strength and hardness of the steel, but its plasticity and toughness clearly decrease. If the carbon content is too high, a clear network of carbides appears in the steel. The presence of network of carbides clearly reduces its strength, plasticity, and toughness, and the strengthening effect of the increased carbon content in the steel is also significantly weakened. Furthermore, it degrades the process performance of the steel, so the carbon content should be kept as low as possible while ensuring strength.

[0057] For duplex steels, the carbon element primarily influences the volume fraction of austenite formed during the annealing process. In the austenite formation process, the diffusion process of carbon elements in austenite or ferrite actually plays a role in controlling the growth process of austenite crystal grains. As the carbon content increases or the heating temperature in the critical region increases, the volume fraction of austenite increases, and furthermore, the martensite phase structure formed after cooling increases, increasing the strength of the material. Considering the synergy between the strength and toughness of the material, and the improvement of strength during the rapid annealing process, the present invention limits the carbon content to the range of 0.04 to 0.12%.

[0058] Manganese (Mn) can form solid solutions with iron, and furthermore, it increases the strength and hardness of ferrite and austenite in carbon steel. In steel materials, relatively fine and high-strength pearlite can be obtained during the cooling process after hot rolling, and the pearlite content can also increase with increasing Mn content. Manganese is also a carbide-forming element, and manganese carbides can dissolve in cementite, thereby indirectly increasing the strength of pearlite. Manganese can also strongly enhance the hardenability of steel, further increasing its strength.

[0059] For duplex steel, manganese is one of the elements that significantly influences the austenite formation dynamics during critical-region annealing. Manganese primarily affects the transformation and growth process to ferrite after austenite formation, as well as the final equilibrium process between austenite and ferrite. Because the diffusion rate of manganese in austenite is much lower than that in ferrite, the growth time of austenite grains controlled by manganese diffusion is long, and the time it takes for manganese to reach a uniform distribution within the austenite is even longer. When heating in the critical region, if the holding time is short, the manganese will not reach a uniform distribution within the austenite, and the subsequent cooling rate will be insufficient, resulting in an inability to obtain a uniform martensite island structure. In duplex steel produced using a rapid heat treatment process (e.g., a water quenching continuous annealing production line), the manganese content is usually high, resulting in a high manganese content after austenite formation, ensuring the hardenability of the austenite islands, and obtaining a relatively uniform performance with a uniform martensite island structure after cooling. Furthermore, since manganese expands the γ phase region and lowers the Ac1 and Ac3 temperatures, manganese-containing steel can obtain a higher martensite volume fraction than low-carbon steel under the same heat treatment conditions. However, high manganese content tends to coarseen the crystal grains in the steel, increasing its susceptibility to overheating; if cooling after molten casting and hot forging / rolling is insufficient, white spots are likely to occur in carbon steel. Taking all of the above into consideration, the present invention is designed with a manganese content in the range of 1.0 to 2.0%.

[0060] Silicon (Si) forms solid solutions in ferrite or austenite, thereby strengthening the yield strength and tensile strength of steel. Furthermore, silicon can increase the cold work deformation hardening rate of steel, making it a beneficial element in alloy steels. In addition, silicon exhibits a clear segregation phenomenon along the crystal cross-section of silicon-manganese steel. The segregation of silicon at the grain boundaries mitigates the distribution of carbon and phosphorus along the grain boundaries, further improving the brittle state of the grain boundaries. Silicon can increase the strength, hardness, and wear resistance of steel without significantly reducing its plasticity. Silicon has a strong deoxidizing ability and is a commonly used deoxidizing agent in steelmaking. Silicon can also increase the fluidity of molten steel. Therefore, silicon is generally contained in steel, but if the silicon content in steel is too high, its plasticity and toughness will be significantly reduced.

[0061] For duplex steel, the main effect of silicon is to reduce the annealing time, which affects the austenite volume fraction at final equilibrium. Although silicon does not have a significant effect on the austenite growth rate, it does have a significant effect on the morphology and distribution of austenite formation. Therefore, in this invention, the silicon content is in the range of 0.1 to 0.5%.

[0062] Cr: The main function of chromium in steel is to increase hardenability, giving the steel relatively good overall mechanical properties after quenching and tempering. Chromium and iron form a continuous solid solution, reducing the austenite phase region, and chromium and carbon form various carbides. Chromium has a greater affinity for carbon than iron and manganese. Chromium and iron can form the intermetallic compound σ phase (FeCr), and chromium reduces the concentration of carbon in pearlite and the critical solubility of carbon in austenite; chromium slows the decomposition rate of austenite, significantly improving the hardenability of steel. However, it also increases the temper brittleness tendency of steel. Chromium increases the strength and hardness of steel, and its effect is more pronounced when used simultaneously with other alloying elements. Cr increases the hardenability of steel during air cooling, thus negatively impacting the weldability of steel. However, if the chromium content is less than 0.3%, the adverse effect on weldability is negligible; above this content, defects such as cracks and slag inclusions are more likely to occur during welding. When chromium (Cr) is present alongside other alloying elements (e.g., with V), its adverse effect on weldability is significantly reduced. For example, when elements such as Cr, Mo, and V are present simultaneously in steel, even if the Cr content reaches 1.7%, there is no significant adverse effect on the weldability of the steel. In this invention, chromium is a beneficial and unnecessary element, and considering factors such as increased cost, too much of it is undesirable. In some embodiments, the Cr content is ≤0.30%.

[0063] Mo (molybdenum) can suppress the self-diffusion of iron and the diffusion rate of other elements. The atomic radius of Mo is larger than that of α-Fe atoms, and when Mo dissolves in an α-solid solution, strong lattice strain is generated in the solution. At the same time, Mo increases the attractive forces of lattice atoms, raising the recrystallization temperature of α-ferrite. The strengthening effect of Mo is evident in pearlite, ferritic, and martensitic steels, and even in high-alloy austenitic steels. The beneficial effect of Mo in steel also depends on its interaction with other alloying elements in the steel. Adding strong carbide-forming elements such as V, Nb, and Ti to steel makes the solid solution strengthening effect of Mo more pronounced. This is because when strong carbide-forming elements combine with C to form stable carbides, more effective dissolution of Mo into the solid solution is promoted, which is advantageous for improving the thermal strength of the steel. The addition of Mo can further improve the hardenability of the steel. Mo suppresses the transformation of the pearlite region and accelerates the transformation in the medium temperature region. Therefore, Mo-containing steel can form a certain number of bainite even at high cooling rates, eliminating the formation of ferrite and pearlite. This is one of the reasons why Mo has a favorable effect on the thermal strength of low-alloy heat-resistant steel. Mo can also significantly reduce the thermal brittleness tendency of steel and decrease the pearlite spheroidization rate. When the Mo content is 0.15% or less, it does not adversely affect the weldability of the steel. In this invention, molybdenum is a beneficial and unnecessary element, and considering factors such as increased cost, too much of it is undesirable. In some embodiments, the Mo content is ≤0.20%.

[0064] Trace alloying elements Ti, Nb, V: When trace amounts of the alloying elements Nb, V, and Ti are added to steel, if the steel has a low carbon equivalent, the dispersion and precipitation of carbide / nitride particles (less than 5 nm in size) and the solid solution of Nb, V, and Ti refine the crystal grains, significantly improving the strength and toughness (especially low-temperature toughness) of the steel and giving it good usability such as weldability. Nb, V, and Ti are carbide and nitride forming elements, and these elements can satisfy these requirements at relatively low concentrations; Nb, V, and Ti are strong carbide forming elements, and at room temperature, they mostly exist in steel in the form of carbides, nitrides, and carbonitrides, with a small amount dissolved in ferrite. Adding Nb, V, and Ti can inhibit the growth of austenite crystal grains and raise the temperature at which the crystal grains of the steel material coarseen. This is because the dispersed carbides and nitrides in the small particles exert a fixing effect on the austenite grain boundaries, inhibiting their movement, raising the austenite recrystallization temperature, and expanding the unrecrystallized region, thus inhibiting the growth of austenite grains. Adding trace amounts of Nb, V, and Ti to steel can, on the one hand, reduce the carbon equivalent content while simultaneously increasing the strength and weldability of the steel, and on the other hand, fix impurities such as oxygen, nitrogen, and sulfur, improving the weldability of the steel; the action of the microparticles, for example, the unsoluble nature of TiN at high temperatures, can prevent grain coarsening in the heat-affected zone, increase the toughness of the heat-affected zone, and thereby improve the weldability of the steel. In this invention, trace alloying elements are beneficial and unnecessary elements, and considering factors such as increased cost, it is undesirable to add too much of them. In some embodiments, the Ti content is ≤0.035%. In some embodiments, the Nb content is ≤0.025%. In some embodiments, the V content is ≤0.030%.

[0065] This invention controls the recovery of the deformed structure, recrystallization, austenite phase, and transformation process in a continuous heat treatment process through a rapid heat treatment process involving rapid heating, short-term heat retention, and rapid cooling. Ultimately, it obtains a fine ferrite structure and a multi-morphological strengthening phase structure, resulting in a material with a good synergy of strength and toughness, reducing alloy costs and the manufacturing difficulty of each process, and improving usability such as welding performance of steel grades with the same strength level.

[0066] The specific principle involves employing different heating rates at different temperature stages of the heating process. In the low-temperature segment, mainly deformation structure recovery occurs, and a relatively low heating rate can be used to reduce energy consumption; in the high-temperature segment, mainly recrystallization of different phase structures and grain growth occur. Unless a relatively high heating rate and short soaking time are used to shorten the residence time of the material in the high-temperature section, it is not possible to ensure that grain growth is small or not large. By controlling the heating rate during the heating process, deformation structure recovery and ferrite recrystallization processes are suppressed, the recrystallization process and the austenite transformation process overlap, increasing the nucleation points of recrystallized grains and austenite grains, and ultimately refining the grains. Short-time holding and rapid cooling shorten the grain growth time during the soaking process, ensuring fine grain structure and uniform distribution.

[0067] The heat treatment process disclosed in Chinese patent application CN10681168B does not segment the entire heating process, and the heating rate employed in the heating process is 20-60°C / s, which is a moderate heating rate. It is realized based on the heating technology of conventional continuous annealing units and cannot be adjusted over a wide range of times as needed for transformation of the material structure.

[0068] In the heat treatment process disclosed in Chinese Patent Application CN107794357B and US Patent Application US2019 / 0153558A1, the heating process is carried out in stages: first, the steel sheet is heated to 300-500°C at a heating rate of 1-10°C / s, then heated to the single-phase austenite region of 850-950°C at a heating rate of 100-500°C / s, and after holding the temperature for no more than 5 seconds, it is cooled to room temperature by water quenching. This processing method requires heating the steel sheet to the high-temperature region of single-phase austenite, which increases the high-temperature resistance requirements of the equipment and increases the difficulty of manufacturing. Furthermore, since a water-cooling method is adopted, the cooling rate is extremely high, which can significantly reduce the growth time of the crystal grain structure in the high-temperature region, but it inevitably leads to non-uniform distribution of alloying elements in the final product, resulting in non-uniform and unstable microstructure performance of the product. Moreover, if the cooling rate is too high, it can cause a series of problems such as defects in the shape of the steel sheet and surface oxidation.

[0069] By comprehensively controlling the entire heat treatment process (including rapid heating (with stage-by-stage heating rate control), short-time soaking, and rapid cooling processes), it is possible to obtain precisely controlled optimal grain size, a uniform distribution of alloying elements and phase structures, and ultimately a product with optimal strength and toughness synergy.

[0070] The average grain size of the ferrite and martensite two-phase structure obtained by the rapid heat treatment method of the present invention is 4 to 10 μm, which is 50 to 60% smaller than the grain size of products produced by conventional techniques (usually 10 to 20 μm). This refinement of the grains increases the strength of the material, and at the same time, good plasticity and toughness are obtained, improving the usability of the material. Furthermore, the ferrite and martensite structures obtained by the present invention have various forms such as massive, strip-like, and granular, and better strength and plasticity can be obtained due to the more uniform distribution of the two.

[0071] The rapid heat treatment method for manufacturing a low-carbon, low-alloy, highly formable duplex steel with a tensile strength of ≥590 MPa, as described in the present invention, comprises the following steps: 1) Smelting and casting The slab is smelted and cast according to the above chemical composition; 2) Hot rolling, coiling Hot rolling end temperature ≥ A r3 Winding temperature: 550-680℃ 3) Cold rolling Cold rolling reduction ratio of 40-85% yields rolled hardened steel strip or steel plate; 4) Rapid heat treatment a) Rapid heating Cold-rolled steel strips or plates are rapidly heated to 750-845°C, with either a one-stage or two-stage rapid heating method being employed; when using a one-stage rapid heating method, the heating rate is 50-500°C / s; when using a two-stage rapid heating method, the first stage involves heating from room temperature to 550-650°C at a heating rate of 15-500°C / s, and the second stage involves heating from 550-650°C to 750-845°C at a heating rate of 30-500°C / s (e.g., 50-500°C / s); b) Soaking The steel strip or sheet is soaked in the austenite-ferrite two-phase region at a target temperature of 750-845°C for a soaking time of 10-60 seconds; c) cooling The uniformly heated steel strip or sheet is slowly cooled to 670-770°C at a cooling rate of 5-15°C / s; then rapidly cooled from 670-770°C to room temperature at a cooling rate of 50-200°C / s; Alternatively, the material is rapidly cooled from 670-770°C to 230-280°C at a cooling rate of 50-200°C / s, followed by overaging treatment, with an overaging treatment time of 200 seconds or less; after overaging treatment, it is cooled to room temperature at a cooling rate of 30-50°C / s.

[0072] Preferably, the total processing time for the rapid heat treatment is 41 to 300 seconds, for example, 41 to 296 seconds.

[0073] Preferably, in step 2), the winding temperature is 580 to 650°C. Preferably, in step 3), the reduction ratio of the cold rolling is 60-80%.

[0074] Preferably, in step 4), when a single-stage heating method is used for the rapid heating, the heating rate is 50 to 300°C / s.

[0075] Preferably, in step 4), a two-stage heating method is used for the rapid heating, in the first stage, heating from room temperature to 550 to 650°C at a heating rate of 15 to 500°C / s, and in the second stage, heating from 550 to 650°C to 750 to 845°C at a heating rate of 50 to 300°C / s.

[0076] Preferably, in step 4), a two-stage heating method is used for the rapid heating, in the first stage, heating from room temperature to 550 to 650°C at a heating rate of 50 to 500°C / s, and in the second stage, heating from 550 to 650°C to 750 to 845°C at a heating rate of 80 to 300°C / s.

[0077] Preferably, in step 4), the final temperature of the rapid heating is 790-830°C.

[0078] Preferably, in step 4), the rapid cooling rate of the steel strip or steel plate is 50 to 150°C / s.

[0079] Preferably, in the soaking process of step 4), the steel strip or steel plate is heated to the target temperature of the austenite and ferrite two-phase region, and then the temperature is maintained to perform soaking.

[0080] Preferably, in the soaking process of step 4), during the soaking time, the steel strip or steel plate is subjected to a small increase or decrease in temperature, so that the temperature after the increase does not exceed 845°C and the temperature after the decrease does not fall below 750°C.

[0081] Preferably, the soaking time is 10 to 40 seconds. Preferably, the overaging treatment time is 20 to 200 seconds.

[0082] Preferably, the duplex steel is a hot-dip galvanized duplex steel described in any one embodiment of this specification, wherein after the soaking step is completed and the steel is slowly cooled to 670-770°C, it is rapidly cooled to 460-470°C at a cooling rate of 50-200°C / s, and then immersed in a zinc pot for hot-dip galvanizing to obtain the hot-dip galvanized duplex steel. In some embodiments, after hot-dip galvanizing, the steel is rapidly cooled to room temperature at a cooling rate of 30-150°C / s to obtain a hot-dip pure galvanized GI product. In some embodiments, after hot-dip galvanizing, the steel is heated to 480-550°C at a heating rate of 30-200°C / s for an alloying treatment, with an alloying treatment time of 10-20 s; after the alloying treatment, the steel is rapidly cooled to room temperature at a cooling rate of 30-250°C / s to obtain an alloyed hot-dip galvanized GA product.

[0083] In some embodiments, the rapid heat treatment method for manufacturing low-carbon, low-alloy, highly formable hot-dip galvanized duplex steel with a tensile strength of ≥590 MPa described herein includes the following steps: A) Smelting and casting The slab is smelted and cast according to the above chemical composition; B) Hot rolling, coiling Hot rolling end temperature ≥ A r3 Winding temperature: 550-680℃ C) Cold rolling After cold rolling with a reduction ratio of 40-85%, rolled hardened steel strip or steel plate is obtained; D) Rapid heat treatment, hot-dip galvanizing a) Rapid heating Cold-rolled steel strips or sheets are rapidly heated from room temperature to a target temperature of 750-845°C in the austenite-ferrite two-phase region, and this rapid heating is carried out in one or two stages; When employing a single-stage rapid heating system, the heating rate is 50-500°C / s; When employing a two-stage rapid heating system, the first stage involves heating from room temperature to 550-650°C at a heating rate of 15-500°C / s, and the second stage involves heating from 550-650°C to 750-845°C at a heating rate of 30-500°C / s (e.g., 50-500°C / s). b) Soaking The austenite-ferrite two-phase region is soaked to a target temperature of 750-845°C for a soaking time of 10-60 seconds; c) Cooling, hot-dip galvanizing After soaking the steel strip or sheet, it is slowly cooled to 670-770°C at a cooling rate of 5-15°C / s; then rapidly cooled to 460-470°C at a cooling rate of 50-150°C / s, and the steel strip or sheet is immersed in a zinc pot for hot-dip galvanizing; d) After hot-dip galvanizing the steel strip or steel plate, rapidly cool it to room temperature at a cooling rate of 50-150°C / s to obtain a hot-dip pure zinc-plated GI product; or, After hot-dip galvanizing of steel strip or steel plate, an alloying treatment is performed by heating to 480-550°C at a heating rate of 30-200°C / s for 10-20 seconds; after the alloying treatment, the material is rapidly cooled to room temperature at a cooling rate of 30-250°C / s to obtain an alloyed hot-dip galvanized GA product.

[0084] Preferably, the total processing time for the rapid heat treatment and hot-dip galvanizing described in step D) is 30 to 142 seconds.

[0085] Preferably, in step B), the winding temperature is 580 to 650°C. Preferably, in step C), the reduction ratio of the cold rolling is 60-80%.

[0086] Preferably, in step D), when a single-stage heating method is used for the rapid heating, the heating rate is 50 to 300°C / s.

[0087] Preferably, in step D), a two-stage heating method is used for the rapid heating, in the first stage, heating from room temperature to 550 to 650°C at a heating rate of 15 to 500°C / s, and in the second stage, heating from 550 to 650°C to 750 to 845°C at a heating rate of 50 to 300°C / s.

[0088] Preferably, in step D), a two-stage heating method is used for the rapid heating, in the first stage, heating from room temperature to 550 to 650°C at a heating rate of 30 to 500°C / s, and in the second stage, heating from 550 to 650°C to 750 to 845°C at a heating rate of 80 to 300°C / s.

[0089] Preferably, in step D), the final temperature of the rapid heating is 770-830°C or 790-830°C.

[0090] Preferably, in the soaking process of step D), the steel strip or steel plate is heated to the target temperature of the austenite and ferrite two-phase region, and then the temperature is maintained to perform soaking.

[0091] Preferably, during the soaking process in step D), the steel strip or steel plate is subjected to a small increase or decrease in temperature during the soaking period, with the temperature after the increase not exceeding 845°C and the temperature after the decrease not falling below 750°C.

[0092] Preferably, the soaking time is 10 to 40 seconds. Preferably, in step D), the steel strip or steel sheet is alloyed and then rapidly cooled to room temperature at a cooling rate of 30-2000°C / s to obtain an alloyed hot-dip galvanized GA product.

[0093] In the rapid heat treatment manufacturing method for 590 MPa class low-carbon low-alloy highly formable duplex steel and highly formable hot-dip galvanized duplex steel described in the present invention: 1. Control of heating rate The recrystallization dynamics of a continuous heating process can be quantitatively described by a relation that is influenced by the heating rate, and the functional relationship between the ferrite recrystallization volume fraction and temperature T in a continuous heating process is:

[0094]

number

[0095] Hereinafter, X(t) is the volume fraction of ferrite recrystallization; n is the Avrami index, which is related to the transformation mechanism and depends on the decay period of the recrystallization nucleation rate, generally taking a value in the range of 1 to 4; T is the heat treatment temperature; T star b(T) is the recrystallization start temperature; β is the heating rate; and b(T) is obtained by the following equation:

[0096]

number

[0097] From the above equations and related experimental data, it can be seen that as the heating rate increases, the recrystallization start temperature (T star ) and end temperature (T fin ) However, both increase; when the heating rate is 50°C / s or higher, the austenite transformation and the recrystallization process overlap, and the recrystallization temperature rises to the two-phase region temperature, and it is known that the faster the heating rate, the higher the ferrite recrystallization temperature.

[0098] In conventional heat treatment processes, due to limitations in heating technology, all heating is slow. Under these conditions, recovery, recrystallization, and grain growth occur sequentially in the deformed matrix, followed by the transformation of ferrite to austenite. Moreover, transformation nucleation mainly occurs at the grown ferrite grain boundaries, and the nucleation rate is relatively low. Therefore, the grain structure of the resulting duplex steel is relatively coarse.

[0099] Under ultra-rapid heating conditions, the deformed matrix begins to undergo transformation from ferrite to austenite just after recrystallization is complete, or even before recrystallization is complete (sometimes not even fully recovered). If recrystallization is complete, or if it is not, the grains are fine and the grain boundary area is large, so the nucleation rate increases significantly, and the austenite grains become clearly finer. At this time, the ferrite recrystallization and austenite transformation processes overlap, and a large number of crystal defects such as dislocations remain in the ferrite crystal, providing a large number of nucleation sites for austenite. Austenite nucleation exhibits explosive nucleation, so the austenite grains become even finer. At the same time, the remaining high-density dislocation line defects also become pathways for the rapid diffusion of carbon atoms, shortening the diffusion time of carbon atoms and allowing each austenite grain to be rapidly generated and grow, so the austenite grains become smaller and the volume fraction increases.

[0100] By precisely controlling the microstructure changes and the distribution of alloying elements and phase components during rapid heating, a good foundation is laid for the growth of the austenite structure during the subsequent soaking process, as well as for the distribution of each alloying component and the transformation of austenite to the martensite phase during the rapid cooling process. Ultimately, a final product microstructure with refined grains and a rational distribution of elements and phases can be obtained. Considering factors such as the effects of rapid heating and refined grains, manufacturing costs, and manufacturability, the present invention sets the heating rate for single-stage rapid heating at 50 to 500°C / s and the heating rate for two-stage rapid heating at 15 to 500°C / s.

[0101] Within different temperature ranges, rapid heating has different effects on microstructural change processes such as material recovery, recrystallization, and grain growth, and therefore, the preferred heating rates also differ for different heating temperature ranges in order to obtain optimal microstructural control: from 20°C to 550-650°C, the heating rate has the greatest effect on the recovery process, and the heating rate should be controlled to 5-300°C / s, more preferably 50-300°C / s; from 550-650°C to the austenitization temperature of 750-845°C, the heating rate has the greatest effect on the nucleation rate and grain growth process, and the heating rate should be controlled to 50-300°C / s, more preferably 80-300°C / s.

[0102] 2. Control of the heating rate The selection of the soaking temperature must be linked to the control of the material's microstructure changes at each temperature stage of the heating process, and at the same time, the changes and control of the microstructure during the subsequent rapid cooling process must be considered, thereby ultimately obtaining a desirable microstructure and distribution.

[0103] Typically, the soaking temperature depends on the carbon content, and the carbon content in the duplex steel of the present invention is 0.04 to 0.12%, and the A of the steel of the present invention C1 and A C3 These are approximately 730°C and 870°C, respectively. The rapid heat treatment process of the present invention raises the soaking temperature of the steel strip to A C1 From A C3 By controlling the temperature within a two-phase region and utilizing rapid heating technology, the material is heated from room temperature to A C1 From A C3 By rapidly heating the material during the process, a large number of dislocations are retained in the ferrite that has not been sufficiently recrystallized, providing a greater nucleation driving force for the transformation of the ferrite into the austenite phase. Compared to traditional continuous annealing processes, the rapid heat treatment method of the present invention can yield a finer austenite structure.

[0104] This invention proposes for the first time that when controlling the soaking temperature, the soaking temperature is raised and lowered within a certain range: that is, a gradient heating and cooling of the soaking temperature is performed, but the soaking temperature must be kept within a certain range. The advantage of doing so is that the process of rapidly raising / lowering the temperature within the temperature range of the two-phase region actually further increases the degree of superheating and supercooling in order to facilitate a rapid transformation process. If the range of temperature increase / decrease is sufficiently large and the rate of temperature increase / decrease is sufficiently large, the repeated transformation of ferrite to the austenite phase and the transformation of austenite to the ferrite phase can further refine the crystal grains. At the same time, it has a certain effect on carbide formation and the uniform distribution of alloying elements, and ultimately it is possible to form alloying elements with a finer structure and uniform distribution.

[0105] Cold-rolled duplex steel contains a large amount of uniformly distributed, undissolved fine carbides. These carbides not only serve as nucleation sites for austenite, but also act as mechanical inhibitors on austenite grain growth during the heating and soaking processes, which is advantageous for refining the grain size of alloy steel. However, if the heating temperature is too high, the number of undissolved carbides decreases significantly, weakening this inhibitory effect, increasing the tendency for grain growth, and further reducing the strength of the steel. If the number of undissolved carbides is too large, it can cause aggregation, leading to a non-uniform distribution of chemical components in the local area. If the carbon content in these aggregated areas is too high, it can cause localized overheating. Therefore, ideally, a small amount of fine, particulate undissolved carbides should be uniformly distributed in the steel. This prevents abnormal growth of austenite grains and, accordingly, increases the content of each alloying element in the matrix, thereby achieving the goal of improving the mechanical properties such as strength and toughness of alloy steel.

[0106] The selection of the soaking temperature must avoid coarsening of the austenite grains in order to obtain fine and uniform austenite grains, thereby achieving the objective of obtaining a fine martensite structure after cooling. A soaking temperature that is too high will coarse the austenite grains, resulting in a coarser martensite structure after rapid cooling, thus reducing the mechanical properties of the steel; it will also increase the number of retained austenites and decrease the number of martensites, reducing the hardness and wear resistance of the steel. A soaking temperature that is too low will not only decrease the number of austenites but also deplete the alloying element content in the austenite, creating an uneven concentration distribution of alloying elements in the austenite, significantly reducing the hardenability of the steel, and negatively impacting its mechanical properties. The soaking temperature for hypoeutectoid steel should be Ac3 + 30~50°C. In ultra-high-strength steel, the soaking temperature can be appropriately increased because the presence of carbide-forming elements inhibits carbide transformation. Taking all of the above factors into consideration, the present invention selects 750-845°C as the soaking temperature and expects to obtain a more ideal and rational final structure.

[0107] 3. Control of soaking time Factors influencing the soaking time also depend on the carbon and alloying element content in the steel. When the carbon and alloying element content in the steel increases, not only does the thermal conductivity of the steel decrease, but the diffusion rate of alloying elements becomes slower than that of carbon elements. Therefore, alloying elements can significantly delay the structural transformation of the steel, and in such cases, the holding time must be appropriately extended. The present invention employs rapid heating, incorporates a large number of dislocations in the two-phase material, provides a large number of nucleation sites for austenite formation, and provides rapid diffusion pathways for carbon atoms, allowing austenite to form very quickly. Moreover, the shorter the soaking and holding time, the shorter the carbon atom diffusion distance, the larger the carbon concentration gradient within the austenite, and finally, the larger the content of remaining retained austenite carbon. However, if the holding time is too short, the distribution of alloying elements in the steel becomes uneven, resulting in insufficient austenitization. If the holding time is too long, the austenite grains tend to become coarse. Therefore, controlling the soaking time must be determined by strictly combining the soaking temperature, rapid cooling, and rapid heating processes and considering them comprehensively; otherwise, it is impossible to ultimately obtain an ideal structure and elemental distribution. For the reasons stated above, the soaking time in this invention is set to 10 to 60 seconds.

[0108] 4. Control of the rapid cooling rate Controlling the rapid cooling process requires combining various factors such as the results of changes in each microstructure during the initial heating and soaking processes, as well as the results of alloy diffusion distribution, ultimately ensuring the acquisition of a material microstructure with ideal phase structures and a rational elemental distribution.

[0109] In order to obtain a martensitic strengthened phase, the cooling rate of the material during rapid cooling must be greater than the critical cooling rate; otherwise, a martensitic structure cannot be obtained. The critical cooling rate mainly depends on the material composition. In this invention, the optimized Si content is 0.1-0.5% and the Mn content is 1.0-2.0%. Since these contents are relatively high, Si and Mn significantly enhance the hardenability of the duplex steel and reduce the critical cooling rate.

[0110] The cooling rate also needs to be determined by comprehensively considering the results of the microstructure changes and alloy diffusion distribution during the heating and soaking processes in order to ultimately obtain a reasonable distribution of each phase structure and alloy element. If the cooling rate is too low, a martensitic structure cannot be obtained, the strength will decrease, and the mechanical performance will not meet the requirements; on the other hand, if the cooling rate is too high, large quenching stresses (i.e., microstructure stresses and thermal stresses) will be generated, causing severe defects in the plate shape, and in particular, if the cooling is uneven, the defects in the plate shape will be especially severe, and it is also prone to severe deformation and cracking of the sample. Therefore, in this invention, the quenching rate is set to 50 to 200°C / s, preferably 50 to 150°C / s.

[0111] 5. Control of over-application Conventional overaging after heat treatment primarily involves tempering the hardened martensite to improve the overall performance of duplex steel. If the overaging temperature and time settings are not appropriate, the martensite will decompose, directly degrading the mechanical performance of the duplex steel. Setting the overaging temperature and time requires comprehensive consideration of the morphology and distribution of the martensite structure, the elemental content and distribution, and the size and distribution of other microstructures. Therefore, controlling overaging requires integrating the parameters of the preceding heating, soaking, and cooling processes. This invention combines the microstructure changes and elemental distribution profiles of rapid heating, short-term holding, and rapid cooling processes, setting the overaging temperature range to 230-280°C and controlling the overaging time to 200 seconds or less.

[0112] 6. Hot-dip galvanizing and alloying control This invention allows for the adjustment and control of parameters in the heat treatment process of steel strips, thereby maintaining better temperature and surface condition during hot-dip galvanizing and alloying treatment, and obtaining superior hot-dip galvanized and alloyed products. By improving the ability to control the surface quality of steel strips, high-surface-quality steel strip products can be obtained.

[0113] This invention modifies traditional continuous annealing units using rapid heating and quenching processes to enable rapid heat treatment. This significantly shortens the length of the heating and soaking segments in conventional continuous annealing furnaces, increasing the production efficiency of conventional continuous annealing units, reducing production costs and energy consumption, decreasing the number of furnace rollers in the continuous annealing furnace, and improving the ability to control the surface quality of the strip steel, thereby achieving high-quality strip steel products. Simultaneously, constructing a new type of continuous annealing unit employing rapid heat treatment process technology simplifies the continuous heat treatment unit, offering advantages such as flexible material transfer and high control capability. For the material, it can refine the grain size of the strip steel, further increasing material strength, lowering alloy costs and the difficulty of manufacturing pre-heat treatment processes, and improving the usability of the material for applications such as welding.

[0114] Based on the above, the present invention, by adopting a rapid heat treatment process, will greatly accelerate the progress of continuous annealing process technology for cold-rolled steel strips. Cold-rolled steel strips can be expected to complete the process from room temperature to the final austenitization within tens of seconds, tens of seconds, or even seconds. This significantly shortens the length of the heating segment of the continuous annealing furnace, improves the speed and production efficiency of the continuous annealing unit, and drastically reduces the number of rollers in the furnace of the continuous annealing unit. In the case of a rapid heat treatment line with a unit speed of about 180 meters / minute, the number of rollers in the high-temperature furnace segment will not exceed 10, and the quality of the steel strip surface can be clearly improved. At the same time, the rapid heat treatment process method for recrystallization and austenitization, which is completed in an extremely short time, also provides a more flexible and soft high-strength steel microstructure design method, improving the material structure and enhancing material performance without changing the pre-process conditions such as alloy composition and rolling process.

[0115] Advanced high-strength steels, such as duplex steel, have broad future potential, and rapid heat treatment technology also has enormous development and application value. The combination of these two technologies provides significant opportunities for the development and production of duplex steel. [Effects of the Invention]

[0116] Compared to the prior art, the present invention has the following advantages: (1) The present invention suppresses the recovery of the deformed structure and the ferrite recrystallization process during the heat treatment process by rapid heat treatment, overlaps the recrystallization process and the austenite transformation process, increases the nucleation points of recrystallized grains and austenite grains, shortens the grain growth time, and results in a two-phase structure of ferrite and martensite in the microstructure of the resulting duplex steel, with an average grain size of 2 to 10 μm, which is significantly smaller than the grain size of products manufactured by conventional techniques (usually 10 to 20 μm); the refinement of the grains increases the strength of the material, and at the same time obtains good plasticity and toughness, improving the usability of the material; in addition, the ferrite and martensite structures obtained in the present invention have various forms such as massive, strip-like, and granular, and the distribution of both is more uniform, resulting in better strength and plasticity.

[0117] (2) Compared to duplex steel obtained by conventional heat treatment methods, the duplex steel obtained by the present invention has a finer grain structure and a more uniform distribution of each phase; therefore, the toughness of the material is clearly improved, and the yield strength of the duplex steel can be controlled within a narrow range of ≥350 MPa, the tensile strength within a narrow range of ≥620 MPa, the elongation within a range of ≥19%, the strength-to-elongation product within a range of ≥15 GPa%, and the strain hardening index n 90 A value greater than 0.2 indicates excellent molding performance.

[0118] (3) The rapid heat treatment process for low-carbon, low-alloy, highly formable duplex steel described in the present invention reduces the total heat treatment time to 41-300 s, and the rapid heat treatment process for low-carbon, low-alloy, highly formable hot-dip galvanized duplex steel reduces the total heat treatment time to 30-142 s, significantly reducing the overall heat treatment process time (traditional continuous annealing process time is usually 5-8 min), thereby increasing production efficiency, reducing energy consumption, and lowering production costs.

[0119] (4) Compared to conventional duplex steel and its heat treatment process, the heating segment and soaking segment times of the rapid heat treatment method of the present invention are reduced by 60-80%, shortening the processing time of the strip steel at high temperatures, saving energy, reducing emissions and consumption, significantly reducing the one-time investment in furnace equipment, and significantly reducing production operating costs and equipment maintenance costs; in addition, by producing products with the same strength level through rapid heat treatment, the alloy content can be reduced, lowering the production costs of heat treatment and pre-processing, and lowering the manufacturing difficulty of each process before heat treatment.

[0120] (5) Compared to duplex steel obtained by conventional continuous annealing, the rapid heat treatment process technology reduces the time of the heating and soaking processes, shortens the furnace length, and reduces the number of furnace rollers, thus reducing the probability of surface defects occurring in the furnace and the probability of defects occurring on the product surface, and significantly improving the surface quality of the product. Furthermore, the refinement of the crystal grains of the product and the reduction in the material alloy content have also improved the usability performance such as forming performance and welding performance of the duplex steel obtained using the technology of the present invention.

[0121] The low-carbon, low-alloy, highly formable duplex steel and hot-dip galvanized duplex steel obtained in this invention have significant value for the development of next-generation lightweight automobiles, trains, ships, aircraft and other means of transportation, as well as the healthy development of corresponding industries and advanced manufacturing sectors. [Brief explanation of the drawing]

[0122] [Figure 1] Figure 1 shows the microstructure image of a duplex steel produced from test steel A according to Example 1 of the present invention. [Figure 2] Figure 2 shows a microstructure image of duplex steel produced from test steel A according to Example 1 of the present invention, following a conventional process 1. [Figure 3] Figure 3 shows the microstructure image of a duplex steel produced from test steel F according to Example 1 of the present invention, in accordance with Example 6. [Figure 4] Figure 4 shows the microstructure image of duplex steel produced from test steel M according to Example 1 of the present invention, in accordance with Example 12. [Figure 5]Figure 5 shows the microstructure image of duplex steel produced from test steel S according to Example 1 of the present invention, in accordance with Example 23. [Figure 6] Figure 6 shows the microstructure image of duplex steel produced from test steel M according to Example 1 of the present invention, in accordance with Example 24. [Figure 7] Figure 7 shows the microstructure image of the duplex steel produced from test steel A according to Example 1 in Example 2 of the present invention. [Figure 8] Figure 8 shows a microstructure image of duplex steel produced from test steel A according to Example 2 of the present invention, following the conventional process 1. [Figure 9] Figure 9 shows a microstructure image of a duplex steel produced from test steel F according to Example 6 of Example 2 of the present invention. [Figure 10] Figure 10 shows the microstructure image of duplex steel produced from test steel M according to Example 2 of the present invention, in accordance with Example 12. [Figure 11] Figure 11 is a microstructure image of a duplex steel produced from test steel S according to Example 2 of the present invention, in accordance with Example 23. [Figure 12] Figure 12 is a microstructure image of duplex steel produced from test steel M according to Example 2 of the present invention, in accordance with Example 24. [Figure 13] Figure 13 shows the microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel A according to Example 1 in Example 3 of the present invention. [Figure 14] Figure 14 shows a microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel A according to Example 3 of the present invention, following the heating rate of the conventional process 1. [Figure 15] Figure 15 shows microstructure images of alloyed hot-dip galvanized duplex steel (GA) produced according to Example 17, from Test Steel I to Example 3 of the present invention. [Figure 16] Figure 16 shows the microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel D according to Example 3 of the present invention, in accordance with Example 22. [Figure 17]Figure 17 shows microstructure images of alloyed hot-dip galvanized duplex steel (GA) produced according to Test Steel I to Example 34 of the present invention. [Figure 18] Figure 18 shows the microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel A according to Example 4 of the present invention, in accordance with Example 1. [Figure 19] Figure 19 shows a microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel A according to Example 4 of the present invention, following the conventional process 1. [Figure 20] Figure 20 shows the microstructure image of alloyed hot-dip galvanized duplex steel (GA) produced from test steel I according to Example 4 of the present invention, in accordance with Example 17. [Figure 21] Figure 21 is a microstructure image of hot-dip pure zinc-plated duplex steel (GI) produced from test steel D according to Example 4 of the present invention, in accordance with Example 22. [Figure 22] Figure 22 shows the microstructure image of alloyed hot-dip galvanized duplex steel (GA) produced from test steel I according to Example 4 of the present invention, in accordance with Example 34. [Modes for carrying out the invention]

[0123] The present invention will be further described below in conjunction with examples and drawings. These examples are based on the technical proposal of the present invention and provide detailed embodiments and specific operating procedures. However, the scope of protection of the present invention is not limited to the following examples.

[0124] In the examples, yield strength, tensile strength, and elongation were determined in accordance with "GB / T228.1-2010 Tensile Tests for Metallic Materials, Part 1: Room Temperature Test Methods," and the test was performed along the transverse direction using sample P7. 90 The test was conducted in accordance with "GB / T228.1-2010 Tensile Testing of Metallic Materials, Part 1: Room Temperature Test Method," using sample P7, along the transverse direction, and based on "GBT 5028-2008 Measurement Method for Extensile Strain Hardening Index (n-value) of Metallic Sheets and Strips," n 90 The value was obtained.

[0125] Example 1 The composition of the test steel of the present invention is shown in Table 1, the specific parameters of this embodiment and the conventional process are shown in Tables 2 and 3, and Tables 4 and 5 show the main properties of the steel produced by the embodiment and the conventional process from the composition of the test steel of this embodiment.

[0126] Tables 1 to 5 show that the rapid heat treatment method of the present invention can reduce the alloy content in the steel to the same level, refine the crystal grains, and obtain a synergy in material structure, strength, and toughness. The yield strength of the duplex steel obtained by the method of the present invention is 350 to 405 MPa, the tensile strength is 624 to 706 MPa, the elongation is 24.4 to 30.4%, the strength-to-elongation product is 17 to 20.1 GPa%, and the strain hardening index is n 90 The value is over 0.20, which is higher than that of duplex steel produced by conventional processes.

[0127] Figure 1 shows the microstructure obtained from typical composition steel A in Example 1, and Figure 2 shows the microstructure obtained from typical composition steel A in Conventional Process Example 1. The images show significant differences between the microstructures treated by different heat treatment methods. The microstructure of steel A treated by the rapid heat treatment process of the embodiment of the present invention (Figure 1) consists of fine grains, mainly a fine martensite structure dispersed on the ferrite matrix, and a small amount of carbides. The microstructure treated by the process of the present invention: ferrite, martensite grain structure, and carbides are all very fine and uniformly distributed in the matrix, which is very advantageous for improving the strength and plasticity of the material. The microstructure of steel A treated by the conventional process (Figure 2) is a typical two-phase steel microstructure image. That is, a small amount of black martensite structure exists at the white ferrite grain boundaries, and due to factors such as elemental segregation, the material microstructure after conventional process treatment exhibits a certain orientation, with the ferrite structure exhibiting a long-length distribution along the rolling direction. The characteristics of microstructures processed using conventional processes are: the grain size is relatively coarse, a certain band-like structure is present, martensite and carbides exhibit a reticular distribution along the ferrite grain boundaries, the ferrite grains are relatively coarse, and the two-phase distribution of ferrite and martensite is not uniform.

[0128] Figure 3 shows the microstructure obtained from typical composition F steel in Example 6 (overaging treatment), and Figure 4 shows the microstructure obtained from typical composition M steel in Example 12 (no aging treatment). Figure 5 shows the microstructure obtained from typical composition S steel in Example 23, and Figure 6 shows the microstructure obtained from typical composition M steel in Example 24. Examples 6, 12, 23, and 24 are all processes with short overall heat treatment cycles. As can be seen from the images, by employing the method of the present invention, even after removing the aging treatment segment, a very uniform, fine, and dispersed phase structure can be obtained. Therefore, the method for manufacturing duplex steel of the present invention can refine the crystal grains, uniformly distribute the phase structure of the material in the matrix, improve the material structure, and enhance material performance.

[0129] [Table 1]

[0130] [Table 2-1]

[0131] [Table 2-2]

[0132] [Table 3-1]

[0133] [Table 3-2]

[0134] [Table 4-1]

[0135] [Table 4-2]

[0136] [Table 5-1]

[0137] [Table 5-2]

[0138] Example 2 Refer to Table 6 for the composition of the test steel in this embodiment, Tables 7 and 8 for the specific parameters of this embodiment and the conventional process, and Tables 9 and 10 show the main properties of the steel produced by this embodiment and the conventional process based on the composition of the test steel in this embodiment.

[0139] Tables 6 to 10 show that the method of the present invention can reduce the alloy content in steel to the same level, refine the crystal grains, and obtain a synergy between the material structure and strength and toughness. The yield strength of the duplex steel obtained by the method of the present invention is 400 to 533 MPa, the tensile strength is 781 to 878 MPa, the elongation is 19.5 to 24.1%, the strength-to-elongation product is 16.3 to 19.3 GPa%, and simultaneously the strain hardening index n 90 The value is over 0.20, which is higher than that of duplex steel produced by conventional processes.

[0140] Figure 7 shows the microstructure obtained from typical composition steel A in Example 1, and Figure 8 shows the microstructure obtained from typical composition steel A in Conventional Process Example 1. The images show significant differences between the microstructures treated by different heat treatment methods. The microstructure of steel A treated by the rapid heat treatment process of the embodiment of the present invention (Figure 7) mainly consists of a fine, uniform martensite structure dispersed on the ferrite matrix and a small amount of carbides. The microstructure treated by the process of the present invention: ferrite, martensite grain structure, and carbides are all very fine and uniformly distributed in the matrix, which is very advantageous for improving the strength and plasticity of the material. The microstructure of steel A treated by the conventional process (Figure 8) is a typical two-phase steel microstructure image. That is, a small amount of black martensite structure exists at the white ferrite grain boundaries, and due to factors such as elemental segregation, the material microstructure after conventional process treatment exhibits a certain orientation, and its ferrite structure exhibits a long-length distribution along the rolling direction. The characteristics of microstructures processed using conventional processes are: coarse grains with the presence of certain band-like structures; martensite and carbides exhibiting a reticular distribution along ferrite grain boundaries; relatively coarse ferrite grains; and a non-uniform distribution of ferrite and martensite two-phase structures.

[0141] Figure 9 shows the microstructure obtained from typical composition F steel in Example 6 (overaging treatment), and Figure 10 shows the microstructure obtained from typical composition M steel in Example 12 (no aging treatment). Figure 11 shows the microstructure obtained from typical composition S steel in Example 23, and Figure 12 shows the microstructure obtained from typical composition M steel in Example 24. Examples 6, 12, 23, and 24 are all processes with short overall heat treatment cycles. As can be seen from the images, by employing the method of the present invention, even after removing the aging treatment segment, a very uniform, fine, and dispersed phase structure can be obtained. Therefore, the method for manufacturing duplex steel of the present invention can refine the crystal grains, uniformly distribute the phase structure of the material in the matrix, improve the material structure, and enhance material performance.

[0142] [Table 6]

[0143] [Table 7-1]

[0144] [Table 7-2]

[0145] [Table 8-1]

[0146] [Table 8-2]

[0147] [Table 9-1]

[0148] [Table 9-2]

[0149] [Table 10-1]

[0150] [Table 10-2]

[0151] Example 3 Refer to Table 11 for the composition of the test steel in this embodiment, and refer to Tables 12 (single-stage heating) and 13 (two-stage heating) for the specific parameters of this embodiment and the conventional process; Tables 14 and 15 show the main properties of GI and GA hot-dip galvanized duplex steels produced by the embodiment and the conventional process shown in Tables 2 and 3, based on the composition of the test steel in this embodiment.

[0152] Tables 11 to 15 show that the method of the present invention can reduce the alloy content in steel to the same level, and obtain a synergy between the material structure and strength and toughness. The yield strength of the duplex steel obtained by the method of the present invention is 347 to 405 MPa, the tensile strength is 624 to 709 MPa, the elongation is 22.2 to 30.3%, the strength-to-elongation product is 15.7 to 19.6 GPa%, and the strain hardening index is n 90 The value is greater than 0.21.

[0153] Figures 13 and 14 show microstructure images of typical steel A after undergoing Example 1 and a comparative conventional process example 1. The two figures reveal significant differences in the microstructure after hot-dip galvanizing. The microstructure of steel A treated with the rapid heat treatment process of the present invention (Figure 1) consists of fine, uniform martensite and carbides dispersed on a fine ferrite matrix. The microstructure treated by the process of the present invention—ferrite, martensite grain structure, and carbides—is extremely fine and uniformly dispersed, which is highly advantageous for improving the material's strength and plasticity. However, the microstructure of steel A treated with a conventional process (Figure 2) is a typical two-phase steel microstructure image. That is, a small amount of black martensite is present at the large white ferrite grain boundaries. The characteristics of the microstructure treated with a conventional process are: coarse grains, a network distribution of martensite and carbides along the ferrite grain boundaries, and a low martensite content.

[0154] Figure 15 shows the microstructure obtained from typical composition I steel in Example 17 (GA), and Figure 16 shows the microstructure obtained from typical composition D steel in Example 22 (GI). Figure 17 shows the microstructure obtained from typical composition I steel in Example 34 (GA). Examples 17, 22, and 34 are all processes with short overall heat treatment cycles. As can be seen from the figures, by employing the rapid heat treatment hot-dip galvanizing method of the present invention, a very uniform, fine, and dispersed phase structure can be obtained even after alloying treatment (Figures 15 and 17). Therefore, the method for manufacturing hot-dip galvanized duplex steel of the present invention can refine the crystal grains, uniformly distribute the phase structure of the material in the matrix, improve the material structure, and enhance material performance.

[0155] [Table 11]

[0156] [Table 12-1]

[0157] [Table 12-2]

[0158] [Table 12-3]

[0159] [Table 13-1]

[0160] [Table 13-2]

[0161] [Table 13-3]

[0162] [Table 14-1]

[0163] [Table 14-2]

[0164] [Table 15-1]

[0165] [Table 15-2]

[0166] Example 4 Refer to Table 16 for the composition of the test steel in this embodiment, and refer to Tables 17 (single-stage heating) and 18 (two-stage heating) for specific parameters of this embodiment and conventional processes; Tables 19 and 20 show the main properties of GI and GA hot-dip galvanized duplex steels produced by the examples and conventional processes in Tables 17 and 18, based on the composition of the test steel in this embodiment.

[0167] Tables 16 to 20 show that the method of the present invention can reduce the alloy content in steel to the same level, and obtain a synergy between the material structure and strength and toughness. The yield strength of the duplex steel obtained by the method of the present invention is 400 to 552 MPa, the tensile strength is 798 to 862 MPa, the elongation is 19.5 to 24.6%, the strength-to-elongation product is 16.3 to 19.9 GPa%, and the strain hardening index is n 90 The value is greater than 0.20.

[0168] Figures 18 and 19 show microstructure images of typical steel A after undergoing Example 1 and a comparative conventional process example 1. The two figures reveal significant differences in the microstructure after hot-dip galvanizing. The microstructure of steel A treated with the rapid heat treatment process of the present invention (Figure 18) consists of fine, uniform martensite and carbides dispersed on a fine ferrite matrix. The microstructure treated by the process of the present invention—ferrite, martensite grain structure, and carbides—is extremely fine and uniformly dispersed, which is highly advantageous for improving the material's strength and plasticity. However, the microstructure of steel A treated with a conventional process (Figure 19) is a typical two-phase steel microstructure image. That is, a small amount of black martensite is present at the grain boundaries of a large white ferrite structure. Due to factors such as elemental segregation, the material microstructure after conventional process treatment exhibits a certain orientation, and its ferrite structure shows a long-length distribution along the rolling direction. The characteristics of microstructures treated with conventional heat treatment processes are as follows: the crystal grains are coarse, and a certain band-like structure exists; martensite and carbides exhibit a network distribution along the ferrite grain boundaries; the ferrite crystal grains are relatively coarse; and the two-phase distribution of ferrite and martensite is not uniform.

[0169] Figure 20 shows the microstructure obtained from typical composition I steel in Example 17 (GA), and Figure 21 shows the microstructure obtained from typical composition D steel in Example 22 (GI). Figure 5 shows the microstructure obtained from typical composition I steel in Example 34 (GA). Examples 17, 22, and 34 are all processes with short overall heat treatment cycles. As can be seen from the figures, by employing the rapid heat treatment hot-dip galvanizing method of the present invention, even after alloying treatment, a very uniform, fine, and dispersed phase structure can be obtained (Figure 22); in the conventional process 9, a coarse ferrite structure is obtained, with a small amount of martensite structure distributed at the ferrite grain boundaries, resulting in a typical hot-dip galvanized duplex steel structure. Therefore, the method for producing hot-dip galvanized duplex steel of the present invention can refine the crystal grains, uniformly distribute the phase structure of the material in the matrix, improve the material structure, and enhance material performance.

[0170] [Table 16]

[0171] [Table 17-1]

[0172] [Table 17-2]

[0173] [Table 17-3]

[0174] [Table 18-1]

[0175] [Table 18-2]

[0176] Table 18-3

[0177] Table 19-1

[0178] Table 19-2

[0179] Table 20-1

[0180] Table 20-2

Claims

1. The chemical composition is characterized by, by mass percentage, C: 0.105-0.12%, Si: 0.1-0.5%, Mn: 1.0-1.16%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02-0.06%, V: 0.025-0.05%, and may further contain one or two from Cr, Mo, Ti, and Nb, with Cr + Mo + Ti + Nb + V ≤ 0.5%, and the remainder being Fe and unavoidable impurities, characterized in a low-carbon low-alloy highly formable duplex steel sheet or a highly formable hot-dip galvanized duplex steel sheet with a tensile strength ≥ 590 MPa. The low-carbon, low-alloy, highly formable duplex steel sheet with a tensile strength of ≥ 590 MPa or the highly formable hot-dip galvanized duplex steel sheet with a tensile strength of ≥ 590 MPa has a yield strength of 340 to 560 MPa, a tensile strength of 620 to 880 MPa, an elongation of 19 to 30.5%, a tensile strength-to-elongation product of 15.5 to 20.5 GPa%, and a strain hardening index n. ９０ The value is greater than 0.20, The microstructure of the low-carbon, low-alloy, highly formable duplex steel sheet with a tensile strength of ≥ 590 MPa or the highly formable hot-dip galvanized duplex steel sheet with a tensile strength of ≥ 590 MPa is a ferrite-martensite duplex structure having an average crystal grain size of 2 to 10 μm.

2. The chemical composition of the highly formable hot-dip galvanized duplex steel sheet having a tensile strength of ≥ 590 MPa is, by mass percentage, C: 0.105 to 0.12%, Si: 0.1 to 0.5%, Mn: 1.0 to 1.16%, P ≤ 0.02%, S ≤ 0.006%, Al: 0.02 to 0.055%, V: 0.025 to 0.05%, and may further contain one or two from Cr, Mo, Ti, and Nb, and Cr + Mo + Ti + Nb + V ≤ 0.5%, with the remainder being Fe and unavoidable impurities, as described in claim 1.

3. The chemical composition of the aforementioned highly formable hot-dip galvanized duplex steel sheet with a tensile strength of ≥ 590 MPa is, by mass percentage, C: 0.105%, Si: 0.1-0.4%, Mn: 1.0-1.16%, P ≤ 0.02%, S ≤ 0.006%, Al: 0.02-0.055%, and may further contain one or two of Cr, Mo, Ti, Nb, and V, with Cr + Mo + Ti + Nb + V ≤ 0.3%, and the remainder being Fe and unavoidable impurities. The yield strength of the aforementioned highly formable hot-dip galvanized duplex steel sheet with a tensile strength of ≥ 590 MPa is 340–410 MPa, the tensile strength is 620–710 MPa, the elongation is 22–30.5%, the tensile strength-to-elongation product is 15.5–20.0 GPa%, and the strain hardening index is n. ９０ The value is greater than 0.21, and the highly formable hot-dip galvanized duplex steel sheet according to claim 1 has a tensile strength of ≥ 590 MPa.

4. A method for manufacturing a low-carbon, low-alloy, highly formable duplex steel sheet having a tensile strength of ≥ 590 MPa as described in claim 1, wherein the low-carbon, low-alloy, highly formable duplex steel sheet having a tensile strength of ≥ 590 MPa has a mass percentage composition of C: 0.105 to 0.12%, Si: 0.1 to 0.5%, Mn: 1.0 to 1.16%, P ≤ 0.02%, S ≤ 0.015%, Al: 0.02 to 0.06%, V: 0.025 to 0.05%, and further containing Cr, M The low-carbon, low-alloy, highly formable duplex steel sheet, which may contain one or two of o, Ti, and Nb, and whose chemical composition is Cr + Mo + Ti + Nb + V ≤ 0.5%, with the remainder being Fe and unavoidable impurities, has a yield strength of 340 to 560 MPa, a tensile strength of 620 to 880 MPa, an elongation of 19 to 30.5%, a tensile strength-to-elongation product of 15.5 to 20.5 GPa%, and a strain hardening index n ９０ The microstructure of the low-carbon, low-alloy, highly formable duplex steel sheet with a value greater than 0.20 and a tensile strength of ≥ 590 MPa is a ferrite-martensite duplex structure having an average grain size of 2 to 10 μm, and the manufacturing method includes the following steps: 1) Smelting and casting The slab is smelted and cast according to the above chemical composition; 2) Hot rolling, coiling Hot rolling end temperature ≥ Ar ３ Winding temperature: 550-680°C; 3) Cold rolling Cold rolling reduction ratio of 40-85% yields rolled hardened steel strip or steel plate; 4) Rapid heat treatment a) Rapid heating Cold-rolled steel strips or plates are rapidly heated to 750-845°C, with either a one-stage or two-stage rapid heating method being employed; when using a one-stage rapid heating method, the heating rate is 50-500°C / s; when using a two-stage rapid heating method, the first stage involves heating from room temperature to 550-650°C at a heating rate of 15-500°C / s, and the second stage involves heating from 550-650°C to 750-845°C at a heating rate of 30-500°C / s; b) Soaking The steel strip or sheet is soaked in the austenite-ferrite two-phase region at a target temperature of 750-845°C for a soaking time of 10-60 seconds; c) cooling The uniformly heated steel strip or sheet is slowly cooled to 670-770°C at a cooling rate of 5-15°C / s; then rapidly cooled from 670-770°C to room temperature at a cooling rate of 50-200°C / s; Alternatively, the material is rapidly cooled from 670-770°C to 230-280°C at a cooling rate of 50-200°C / s, followed by overaging treatment, with an overaging treatment time of 200 s or less; after overaging treatment, it is cooled to room temperature at a cooling rate of 30-50°C / s.

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