CN115181898A

CN115181898A - 1280 MPa-grade low-carbon low-alloy Q & P steel and rapid thermal treatment manufacturing method thereof

Info

Publication number: CN115181898A
Application number: CN202110360562.7A
Authority: CN
Inventors: 李俊; 刘赓; 王健; 王超; 陈云鹏; 刘益民
Original assignee: Baoshan Iron and Steel Co Ltd
Current assignee: Baoshan Iron and Steel Co Ltd
Priority date: 2021-04-02
Filing date: 2021-04-02
Publication date: 2022-10-14
Anticipated expiration: 2041-04-02
Also published as: CN115181898B

Abstract

A1280 MPa-level low-carbon low-alloy Q & P steel and a rapid heat treatment manufacturing method thereof are disclosed, wherein the steel comprises the following chemical components in percentage by mass: 0.16 to 0.23 percent of C, 1.4 to 2.0 percent of Si, 2.4 to 3.0 percent of Mn, 0.006 to 0.016 percent of Ti, less than or equal to 0.015 percent of P, less than or equal to 0.002 percent of S, 0.02 to 0.05 percent of Al, and one or two of Cr, mo, nb and V, less than or equal to 0.5 percent of Cr + Mo + Ti + Nb + V, and the balance of Fe and other inevitable impurities. The manufacturing method comprises the following steps: smelting, casting, hot rolling, cold rolling and rapid heat treatment; the whole process of the rapid thermal treatment takes 71 to 186s. The invention changes the recovery, recrystallization and austenite phase transformation processes of the deformed structure, increases the nucleation rate (including the recrystallization nucleation rate and the austenite phase transformation nucleation rate), shortens the grain growth time, refines the grains, improves the mechanical property of the material while improving the heat treatment efficiency and expands the range of the material property interval by controlling the rapid heating, short-time heat preservation and rapid cooling processes in the rapid heat treatment process.

Description

1280 MPa-grade low-carbon low-alloy Q & P steel and rapid heat treatment manufacturing method thereof

Technical Field

The invention belongs to the technical field of rapid heat treatment of materials, and particularly relates to 1280 MPa-grade low-carbon low-alloy Q & P steel and a rapid heat treatment manufacturing method thereof.

Background

With the increasing awareness of energy conservation and material safety service, the use of high-strength steel, especially advanced high-strength steel, is increasing, which makes the development of advanced high-strength steel more and more important for iron and steel enterprises and scientific research institutes. In order to further increase the product of strength and elongation of steel products, the development of third generation advanced high strength steels represented by Q & P (Quenching and arresting, quenching and carbon redistribution) steels has been increasingly paid attention.

The Q & P heat treatment process is a novel continuous heat treatment process technology provided by Speer et al in the beginning of the 21 st century, and mainly comprises four steps:

firstly, heating the strip steel to austenitizing temperature and preserving heat;

second, rapidly cool the sample to M _s ～M _f At a certain temperature to obtain a dual-phase structure mainly comprising martensite and retained austenite;

thirdly, heating the strip steel to be not higher than M _s The temperature is kept, so that carbon elements are diffused and distributed from supersaturated martensite to austenite, the carbon content and hardness of the martensite are reduced, the plasticity of the martensite is improved, and the carbon content of the austenite is improved and the stability of the austenite is improved;

and fourthly, cooling to room temperature, wherein if the stability of the retained austenite is insufficient in the process, part of the austenite is transformed into martensite, and the amount of the retained austenite obtained at room temperature is reduced.

The Q & P steel is essentially a martensitic steel, but is different from the traditional tempered martensitic steel, and the plasticity of the Q & P steel is greatly improved under the same strength with the tempered martensitic steel. This is because the structure of Q & P steel contains retained austenite, which is transformed into martensite during the deformation process, and the so-called TRIP effect is generated, thereby greatly improving the plasticity of the steel.

At present, two development means aiming at the Q & P process are provided, and firstly, the inhibition capacity of alloy elements in steel on carbide precipitation is improved by adding the alloy elements; and secondly, optimizing the process, namely changing the structure performance of the Q & P steel by adjusting the temperature and time of the quenching and distribution process in the Q & P process.

U.S. patent No. 2003/027825 proposes a general process of Q & P steel production, and limits the austenitizing process to be performed at high temperature, the material structure needs to be fully austenitized, the temperature for the actual production process is too high (850-950 ℃), the time is long (usually, the steel plate austenitizing process needs to be kept for 2-5 min), the equipment requirement is high, and the manufacturing cost is high.

Chinese patent CN103667884B discloses a method for preparing 1400MPa grade cold-rolled ultrahigh strength automobile steel with low yield ratio and high elongation, the steel of the invention comprises the following components by mass percent: 0.14 to 0.16 percent of C, 1.31 to 1.51 percent of Si, 2.7 to 2.9 percent of Mn, less than or equal to 0.005 percent of S, less than or equal to 0.009 percent of P, 0.11 to 0.51 percent of Al, 0.005 to 0.02 percent of RE, and the balance of Fe and other inevitable impurity elements. The steel is mainly characterized in that high Mn and rare earth elements are added to reduce Mn element segregation, a certain amount of metastable austenite structure is obtained through a traditional continuous annealing process, the tensile strength is greater than 1400MPa, the yield strength is 500-900MPa, the elongation is greater than 8%, and the yield ratio is 0.4-0.6.

Chinese patent CN102925799A discloses a production method of ultra-high strength steel, and the chemical components of the steel in percentage by mass are as follows: 0.21 to 0.24 percent of C, 0.45 to 0.55 percent of Si, 1.4 to 1.46 percent of Mn, less than or equal to 0.015 percent of P, less than or equal to 0.01 percent of S, 0.03 to 0.06 percent of Al, 0.02 to 0.03 percent of Nb, 0.05 to 0.06 percent of V, 0.035 to 0.05 percent of Ti, 0.25 to 0.35 percent of Mo, 0.0017 to 0.0022 percent of B, and the balance of Fe and other inevitable impurities. The yield strength of the obtained steel plate was about 1300MPa, the tensile strength was about 1378MPa, and the elongation after fracture was about 11%. The steel is mainly characterized in that high alloy content is added, including high C element and high micro alloy element, and Mo element is added. It is processed by a conventional continuous annealing process, and the annealing temperature is as high as 910-930 ℃, which has high requirements on manufacturing cost and manufacturing equipment.

Chinese patent CN105543674B discloses a method for manufacturing cold-rolled ultrahigh-strength dual-phase steel with high local forming performance, and the chemical components of the high-strength dual-phase steel comprise the following components in percentage by weight: c:0.08 to 0.12%, si:0.1 to 0.5%, mn:1.5 to 2.5%, al:0.015 to 0.05 percent, and the balance of Fe and other inevitable impurities. Selecting and matching raw materials for the chemical components, and smelting into a casting blank; heating the casting blank at 1150-1250 ℃ for 1.5-2 hours, and then carrying out hot rolling, wherein the initial rolling temperature of the hot rolling is 1080-1150 ℃, and the final rolling temperature is 880-930 ℃; cooling to 450-620 ℃ at a cooling speed of 50-200 ℃/s after rolling, and coiling to obtain a hot rolled steel plate with bainite as a main structure type; and (3) cold rolling the hot rolled steel plate, heating to 740-820 ℃ at the speed of 50-300 ℃/s, annealing, keeping the temperature for 30s-3min, cooling to 620-680 ℃ at the cooling speed of 2-6 ℃/s, and then cooling to 250-350 ℃ at the cooling speed of 30-100 ℃/s for overaging treatment for 3-5min to obtain the ferrite-martensite dual-phase structure ultrahigh-strength dual-phase steel. The yield strength of the ultrahigh-strength dual-phase steel is 650-680MPa, the tensile strength is 1023-1100MPa, the elongation is 12.3-13%, and the ultrahigh-strength dual-phase steel does not crack when bent at 180 degrees along the rolling direction.

The method is mainly characterized in that the control of cooling conditions after hot rolling is combined with the rapid heating in the continuous annealing process, namely, the cooling process after hot rolling is controlled to eliminate banded structures and realize the homogenization of the structures; and rapid heating is adopted in the subsequent continuous annealing process, so that the tissue thinning is realized on the basis of ensuring the tissue uniformity. Therefore, the patent technology adopts rapid heating annealing, and the premise is that the hot rolling raw material with bainite as a main structure is obtained after hot rolling, and the purpose is mainly to ensure the uniformity of the structure and avoid the defect of local deformation caused by the occurrence of banded structures.

The defects of the patent mainly lie in that:

firstly, the hot rolling raw material with bainite structure is obtained, and has high strength and large deformation resistance, thereby bringing great difficulty to subsequent pickling and cold rolling production;

secondly, the understanding of the rapid heating is limited to shortening the heating time and refining the layer of crystal grains, the heating rate is not divided according to the change of the material structure of different temperature sections, and the material is heated at the speed of 50-300 ℃/s, so that the production cost of the rapid heating is increased;

thirdly, the soaking time is 30s-3min, and the increase of the soaking time inevitably and partially weakens the grain refining effect generated by rapid heating, thus being unfavorable for improving the strength and the toughness of the material;

fourth, the patent must be overaged for 3-5 minutes, which is actually too long for rapid heat treating DP steels and is not necessary. And the increase of soaking time and overaging time is not favorable for saving energy, reducing the investment of unit equipment and the occupied area of the unit, and is also not favorable for the high-speed stable operation of the strip steel in the furnace, which is obviously not a rapid heat treatment process in a strict sense.

Chinese patent 201711385126.5 discloses 'a 780 MPa-grade low-carbon low-alloy TRIP steel', which comprises the following chemical components in percentage by mass: c:0.16-0.22%, si:1.2-1.6%, mn:1.6-2.2%, the balance being Fe and other unavoidable impurity elements, obtained by a rapid thermal treatment process comprising: rapidly heating the strip steel to a two-phase region of austenite and ferrite at the temperature of 790-830 ℃ from room temperature, wherein the heating rate is 40-300 ℃/s; heating the target temperature in the two-phase region for 60-100 s; rapidly cooling the strip steel from the temperature of the two-phase region to 410-430 ℃ at the cooling speed of 40-100 ℃/s, and staying in the temperature region for 200-300s; and rapidly cooling the strip steel from the temperature of 410-430 ℃ to the room temperature. The method is characterized in that: the TRIP steel metallographic structure is a bainite, ferrite and austenite three-phase structure; the average grain size of the TRIP steel is obviously refined; the tensile strength is 950-1050 MPa; the elongation is 21-24%; the maximum product of strength and elongation can reach 24GPa%.

The defects of the patent mainly comprise the following aspects:

firstly, the patent discloses a 780 MPa-grade low-carbon low-alloy TRIP steel product and a process technology thereof, but the tensile strength of the TRIP steel product is 950-1050 MPa, the tensile strength of the TRIP steel product is too high as the tensile strength of the 780 MPa-grade product, the use effect of a user is not good, and the tensile strength of the TRIP steel product is too low as the tensile strength of the 980 MPa-grade product, so that the strength requirement of the user can not be well met;

secondly, the patent adopts one-stage rapid heating, the same rapid heating rate is adopted in the whole heating temperature interval, the materials are not processed differently according to the change of the organizational structures of the materials in different temperature sections, and the materials are all rapidly heated at the speed of 40-300 ℃/s, so that the production cost of the rapid heating process is inevitably increased;

thirdly, the soaking time of the patent is set to be 60-100s, which is similar to that of the traditional continuous annealing, and the increase of the soaking time inevitably partially weakens the grain refining effect generated by rapid heating, thus being very unfavorable for improving the strength and the toughness of the material;

fourth, the patent must perform a bainite isothermal treatment time of 200-300s, which is actually too long for rapid heat treatment of the product to function as intended and is not necessary. And the increase of soaking time and isothermal treatment time is not beneficial to saving energy, reducing unit equipment investment and unit occupied area, and is also not beneficial to the high-speed stable operation of the strip steel in the furnace, obviously, the rapid heat treatment process is not in strict meaning.

Chinese patent CN107794357B and US2019/0153558A1 disclose 'a method for producing an ultra-high strength martensite cold-rolled steel sheet by an ultra-fast heating process', wherein the chemical components of the high strength dual-phase steel are as follows by weight percent: c:0.10 to 0.30%, mn:0.5 to 2.5%, si:0.05 to 0.3%, mo:0.05 to 0.3%, ti:0.01 to 0.04%, cr:0.10 to 0.3%, B:0.001 to 0.004 percent, less than or equal to 0.02 percent of P, less than or equal to 0.02 percent of S, and the balance of Fe and other inevitable impurities. The mechanical properties of the dual-phase steel are as follows: yield strength Rp _0.2 Greater than 1100MPa, tensile strength R _m =1800-2300MPa, the maximum elongation of 12.3% and the uniform elongation of 5.5-6%. The invention provides a super-fast heating production process of a super-strength martensite cold-rolled steel plate, which is characterized in that the cold-rolled steel plate is heated to 300-500 ℃ at the speed of 1-10 ℃/s, and then is heated to a single-phase austenite region 850-950 ℃ at the heating rate of 100-500 ℃/s; and then, immediately cooling the steel plate to room temperature after keeping the temperature for not more than 5 seconds to obtain the ultrahigh-strength cold-rolled steel plate.

The disadvantages of the process described in this patent include:

firstly, the annealing temperature of the steel enters the ultra-high temperature range of an austenite single-phase region, and the steel also contains more alloy elements, and the yield strength and the tensile strength both exceed 1000MPa, so that the steel brings great difficulty to the manufacturing of the heat treatment process and the prior process of the heat treatment and the use of subsequent users;

secondly, the ultra-fast heating annealing method adopts the heat preservation time not more than 5s, so that the controllability of the heating temperature is poor, and the distribution of alloy elements in a final product is uneven, so that the structural performance of the product is uneven and unstable;

thirdly, the final rapid cooling adopts water quenching to cool to room temperature without necessary tempering treatment, so that the structure property of the obtained final product and the distribution condition of alloy elements in the final structure can not ensure that the product can obtain the optimal obdurability, the final product has excessive strength and insufficient plasticity and toughness;

fourthly, the method of the invention causes the problems of poor plate shape, surface oxidation and the like of the steel plate due to overhigh water quenching speed, so the patented technology has no high or low practical application value.

At present, limited by the equipment capacity of the traditional continuous annealing furnace production line, the cold-rolled Q & P steel products and related researches of the annealing process are based on the heating rate (5-20 ℃/s) of the existing industrial equipment to slowly heat the strip steel so as to sequentially complete recrystallization and austenitizing phase change, so that the heating time and the soaking time are long, the energy consumption is high, meanwhile, the traditional continuous annealing production line has the defects that the number of rollers of the strip steel in a high-temperature furnace section is large, the traditional continuous annealing unit generally requires 1-3 min according to the requirements of product outline and capacity, the number of rollers in the high-temperature furnace section is generally different from 20-40 for the traditional production line with the unit speed of 180 m/min, and the difficulty in controlling the surface quality of the strip steel is increased.

Disclosure of Invention

The invention aims to provide 1280 MPa-level low-carbon low-alloy Q & P steel and a rapid heat treatment manufacturing method thereof, which change the recovery, recrystallization and austenite phase transformation processes of a deformed structure through rapid heat treatment, increase the nucleation rate (including the recrystallization nucleation rate and the austenite phase transformation nucleation rate), shorten the grain growth time, refine grains, improve the content of residual austenite, and further improve the strength and the plasticity of materials; the yield strength is 754-1112 MPa, the tensile strength is 1235-1350 MPa, the elongation is 19-22.2%, and the product of strength and elongation is 24.8-28.97 GPa%, so that the material can obtain good toughness matching, and the use performance of users such as forming, welding and the like of the material is improved; meanwhile, the rapid heat treatment process is adopted, so that the production efficiency is improved, and the alloy content in the steel of the same grade is reduced, thereby reducing the production cost and the manufacturing difficulty of the working procedure before heat treatment, obviously reducing the number of furnace rollers, and improving the surface quality.

In order to achieve the purpose, the technical scheme of the invention is as follows:

1280 MPa-grade low-carbon low-alloy Q & P steel comprises the following chemical components in percentage by mass: c:0.16 to 0.23%, si:1.4 to 2.0%, mn:2.4 to 3.0%, ti: 0.006-0.016%, P is less than or equal to 0.015%, S is less than or equal to 0.002%, al: 0.02-0.05%, and one or two of Cr, mo, nb and V, wherein Cr + Mo + Ti + Nb + V is less than or equal to 0.5%, and the balance of Fe and other inevitable impurities, and is obtained by the following process:

1) Smelting and casting

Smelting according to the chemical components and casting into a plate blank;

2) Hot rolling and coiling

The coiling temperature is 550-680 ℃;

3) Cold rolling of steel

The cold rolling reduction rate is 40-85%;

4) Rapid thermal processing

Rapidly heating the cold-rolled steel plate to 770-845 ℃, wherein the rapid heating adopts a one-section type or two-section type; when one-section type rapid heating is adopted, the heating rate is 50-500 ℃/s, when two-section type rapid heating is adopted, the first section is heated from room temperature to 550-625 ℃ at the heating rate of 15-500 ℃/s, and the second section is heated from 550-625 ℃ to 770-845 ℃ at the heating rate of 50-500 ℃/s; then soaking for 10-60 s at 770-845 ℃;

after the heat equalization is finished, slowly cooling to 700-770 ℃ at a cooling rate of 5-15 ℃/s, then rapidly cooling to 230-280 ℃ at a cooling rate of 50-200 ℃/s, preserving heat for 2-10 s in the temperature range, then heating to 300-470 ℃ at a heating rate of 10-30 ℃/s for tempering treatment, wherein the tempering time is 10-60 s; after the tempering is finished, cooling to room temperature at a cooling rate of 30-100 ℃/s.

Preferably, the content of C is 0.18 to 0.21%.

Preferably, the Si content is 1.6 to 1.8%.

Preferably, the Mn content is 2.6 to 2.8%.

Preferably, in the step 2), the hot rolling temperature is more than or equal to A _r3 。

Preferably, in the step 2), the coiling temperature is 580 to 650 ℃.

Preferably, in the step 3), the cold rolling reduction is 60 to 80%.

Preferably, the time used in the whole rapid thermal treatment process is 71 to 186s.

Preferably, the rapid heating in the step 1) adopts one-stage heating, and the heating rate is 50-300 ℃/s.

Preferably, the rapid heating in step 1) adopts two-stage heating: the first section is heated from room temperature to 550-625 ℃ at the heating rate of 15-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 50-300 ℃/s.

Preferably, the rapid heating in step 1) adopts two-stage heating: the first section is heated from room temperature to 550-625 ℃ at the heating rate of 30-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 80-300 ℃/s.

Preferably, in the step 4), the steel strip or the steel plate is rapidly cooled from 700-770 ℃ to 230-280 ℃ at a cooling rate of 50-150 ℃/s.

The metallographic structure of the Q & P steel is a multiphase structure comprising 80-90% of martensite, 10-20% of retained austenite and 3-5% of ferrite, the matrix structure is uniformly distributed, obvious lamellar tempered martensite appears, the grain diameter is 1-3 mu m, uniformly distributed ferrite phase exists around martensite strengthening phase grains, and the martensite strengthening phase grains are mainly of a lamellar structure.

The austenite in the metallographic structure of the Q & P steel has good thermal stability, the austenite transformation rate at-50 ℃ is lower than 8%, and the austenite transformation rate at-190 ℃ is lower than 30%.

The Q & P steel has the yield strength of 754-1112 MPa, the tensile strength of 1281-1350 MPa, the elongation of 19-22.2 percent and the product of strength and elongation of 24.8-28.97 GPa percent.

In the composition and process design of the steel of the invention:

c: carbon is the most common strengthening element in steel, and increases the strength and reduces the plasticity of steel, but for forming steel, low yield strength, high uniform elongation and total elongation are required, so the carbon content is not too high. There are two ways that carbon phases exist in steel: ferrite and cementite. The carbon content has great influence on the mechanical properties of the steel, the quantity of strengthening phases such as martensite, pearlite and the like can be increased along with the increase of the carbon content, so that the strength and the hardness of the steel are greatly improved, but the plasticity and the toughness of the steel can be obviously reduced, if the carbon content is too high, obvious net-shaped carbides can appear in the steel, the strength, the plasticity and the toughness of the steel can be obviously reduced due to the existence of the net-shaped carbides, the strengthening effect generated by the increase of the carbon content in the steel can be obviously weakened, the technological properties of the steel are poor, and the carbon content is reduced as much as possible on the premise of ensuring the strength.

For Q & P steel, carbon is one of the most effective strengthening elements of martensite matrix, and is dissolved in austenite to enlarge austenite phase region, greatly raise austenite stability, shift the transformation C curve of pearlite and bainite to the right, delay the transformation of pearlite and bainite and lower the Ms point temperature. Too low a carbon content reduces the stability of the retained austenite, and too high a carbon content causes twins in the martensite, reducing the plasticity, toughness and weldability of the steel. The carbon content is limited to the range of 0.16-0.23% by comprehensively considering the patent.

Mn: manganese can form a solid solution with iron, so that the strength and hardness of ferrite and austenite in the carbon steel are improved, fine pearlite with higher strength is obtained in the cooling process of the steel after hot rolling, and the content of the pearlite is increased along with the increase of the content of Mn. Manganese is a forming element of carbide at the same time, and the carbide of manganese can be dissolved into a cementite, so that the strength of strengthening phases such as martensite and pearlite is indirectly enhanced. Manganese also strongly enhances the hardenability of the steel, further improving its strength.

For Q & P steel, the addition of manganese element can reduce the martensite transformation temperature Ms, increase the content of retained austenite and improve the stability of the retained austenite, and the influence of the manganese element on the toughness of the steel is not large. However, when the manganese content is high, the steel tends to coarsen crystal grains, and the steel is increased in the sensitivity to overheating, and white spots are easily generated in the carbon steel when cooling is not proper after melt casting and hot rolling. The invention limits the manganese content to be within the range of 2.4-3.0%.

Si: silicon forms a solid solution in ferrite or austenite, thereby enhancing the yield strength and tensile strength of the steel, and silicon increases the cold working deformation hardening rate of the steel. In addition, silicon has an obvious enrichment phenomenon on the surface of a fracture along the grain boundary of the silicon-manganese steel, and the segregation of silicon at the position of the grain boundary can slow down the distribution of carbon and phosphorus along the grain boundary, so that the embrittlement state of the grain boundary is improved. Silicon can improve the strength, hardness and wear resistance of steel, and can not obviously reduce the plasticity of the steel within a certain range. Silicon has strong deoxidizing capacity, is a common deoxidizing agent in steel making, and generally contains silicon because the silicon can also increase the fluidity of molten steel, but when the content of the silicon in the steel is too high, the plasticity and the toughness of the steel are obviously reduced. For Q & P steels:

first, si is a non-carbide-forming element, has extremely low solubility in carbide, and suppresses Fe in the isothermal process of QP steel ₃ C is formed to enrich carbon in the untransformed austenite, thereby greatly improving the stability of the austenite and enabling the austenite to be reserved at room temperature;

secondly, the silicon element is a ferrite forming element, so that the stability of the residual austenite can be improved, and the solid solution strengthening effect is achieved, so that the strength of the steel is improved;

thirdly, the silicon element has the function of reducing the austenite phase region and improving the activity of the C element in the ferrite.

A higher silicon content is advantageous for obtaining more retained austenite, but too high a silicon content may cause a hard oxide layer, poor surface properties, and a decrease in wettability and surface quality of the hot rolled steel sheet. Silicon has no obvious influence on the growth rate of austenite, but has obvious influence on the shape and distribution of the austenite, and the increase of the silicon content increases the manufacturing difficulty of the working procedure before heat treatment; the present invention limits the silicon content to a range of 1.4 to 2.0%.

Ti: ti is a microalloy element, belongs to a ferrite forming element of a closed gamma region, can improve the critical point of steel, and Ti and C in the steel can form very stable TiC which is extremely difficult to dissolve in the austenitizing temperature range of general heat treatment. Since TiC particles refine austenite grains, the chances of new phase nucleation increase during austenite decomposition transformation, which accelerates austenite transformation. Ti forms TiC and TiN precipitates with C and N, and is more stable than carbonitride of Nb and V, and significantly reduces the diffusion rate of C in austenite to significantly reduce the austenite formation rate, and the formed carbonitride precipitates in the matrix and pins at the austenite grain boundary to inhibit the austenite grain growth. In the cooling process, the precipitated TiC has the precipitation strengthening effect; in the tempering process, ti slows down the diffusion of C in an alpha phase, slows down the precipitation and growth of carbides such as Fe, mn and the like, increases the tempering stability, and can play a role in secondary hardening through TiC precipitation. The high temperature strength of the steel can be improved by microalloying of Ti. By adding a trace amount of Ti into the steel, on one hand, the strength and the welding performance of the steel can be improved while the carbon equivalent content is reduced; on the other hand, impurities such as oxygen, nitrogen, sulfur, etc. are fixed, thereby improving weldability of steel; secondly, due to the effect of microscopic particles, such as insolubility of TiN at high temperature, coarsening of grains in the heat affected zone is prevented, toughness of the heat affected zone is improved, and thus weldability of steel is improved. In the invention, ti is a beneficial and necessary additive element, and the addition amount is not too much considering factors such as cost increase, and the like, and is preferably 0.006-0.016%.

Cr: the main function of chromium in steel is to improve hardenability, so that the steel has better comprehensive mechanical properties after quenching and tempering. Chromium forms a continuous solid solution with iron, narrowing the austenite phase region, forms multiple carbides with carbon, and has a greater affinity for carbon than the elements iron and manganese. Chromium and iron may form an intermetallic sigma phase (FeCr), chromium reducing the concentration of carbon in pearlite and the limiting solubility of carbon in austenite; chromium slows down the decomposition speed of austenite and obviously improves the hardenability of steel. But also increases the temper brittleness tendency of the steel. When other alloy elements are added, the chromium element has obvious effect of improving the strength and the hardness of the steel. Since Cr increases the quenching ability of steel during air cooling, it adversely affects the weldability of steel. However, when the chromium content is less than 0.3%, the adverse effect on weldability is negligible; if the content is more than this, defects such as cracks and slag inclusion are likely to occur during welding. When Cr is present with other alloying elements (e.g., with V), the adverse effect of Cr on weldability is greatly reduced. If Cr, mo, V, etc. are present in the steel at the same time, the weld properties of the steel are not significantly adversely affected even if the Cr content reaches 1.7%. The chromium element is a beneficial and unnecessary addition element, and the addition amount is not suitable to be too much in consideration of factors such as cost increase and the like.

Mo: molybdenum can inhibit the self-diffusion of iron and the diffusion speed of other elements. The atomic radius of Mo is larger than that of alpha-Fe atoms, so that when Mo is dissolved in the alpha solid solution, the solid solution generates strong lattice distortion, and meanwhile, the crystal lattice atomic bond attraction can be increased by Mo, and the recrystallization temperature of alpha ferrite is increased. The strengthening effect of Mo in pearlite type, ferrite type and martensite type steel and even in high-alloy austenitic steel is also very obvious. The good effect of Mo in steel also depends on the interaction with other alloying elements in the steel. When strong carbide forming elements V, nb and Ti are added into steel, the solid solution strengthening effect of Mo is more obvious. This is because, when a strong carbide-forming element is combined with C into a stable carbide, mo can be promoted to be more efficiently dissolved into solid solution, thereby contributing more to the improvement of the heat strength of the steel. Addition of Mo can also increase the hardenability of the steel, but the effect is less pronounced than C and Cr. Mo inhibits the transformation of pearlite areas and accelerates the transformation in the intermediate temperature area, so that Mo-containing steel can form a certain amount of bainite under the condition of a high cooling speed, and the formation of ferrite is eliminated, which is one of the reasons why Mo favorably influences the heat strength of low-alloy heat-resistant steel. Mo also significantly reduces the hot embrittlement tendency of the steel and reduces the pearlite nodularisation rate. When the Mo content is 0.15% or less, the weldability of the steel is not adversely affected. The molybdenum element is a beneficial and unnecessary addition element, and the addition amount is not too large in consideration of factors such as cost increase and the like.

Nb: nb element is a carbide and nitride forming element and can satisfy such a requirement at a relatively low concentration. At normal temperature, most of the steel exists in the form of carbide, nitride, or carbonitride, and a small amount thereof is dissolved in the ferrite. The addition of Nb can prevent austenite grains from growing and improve the coarsening temperature of steel grains. The Nb element and carbon generate stable NbC, and the addition of trace Nb element in steel can improve the strength of matrix by utilizing the precipitation strengthening effect. The Nb element has obvious inhibition effect on the growth of ferrite recrystallization and the growth of austenite grains, can refine the grains and improve the strength and toughness of the steel; nb can influence the mobility of grain boundaries, and also has an effect on the phase transition behavior and the formation of carbides. Nb can increase the content of carbon in the residual austenite, hinder the formation of bainite, promote martensite nucleation, obtain a dispersed martensite structure, improve the stability of the residual austenite, improve the strength of the dual-phase steel by adding Nb element, obtain the dual-phase steel with certain strength under the conditions of low content of martensite and low content of C, and improve the obdurability of the dual-phase steel; an additional benefit of the simultaneous addition of Nb is that the strength of the steel can be increased over a wider annealing temperature range. In the invention, the Nb element is a beneficial and unnecessary additive element, and the addition amount is not suitable to be excessive in consideration of factors such as cost increase and the like.

Microalloy element V: v is a ferrite stabilizing element and a strong carbide forming element, has strong grain refining effect and can make the steel structure compact. The addition of V to steel can improve the strength, plasticity and toughness of steel at the same time. Vanadium also improves the high temperature strength of structural steels. Vanadium does not improve hardenability. The addition of trace microalloy element V in the steel can ensure that the steel has good weldability and other service performances by the dispersion precipitation of carbon and nitride particles (the size is less than 5 nm) and the solid solution of V to refine grains under the condition of low carbon equivalent. A trace amount of V is added into the steel, so that on one hand, the strength and the welding performance of the steel can be improved while the carbon equivalent content is reduced; on the other hand, impurities such as oxygen, nitrogen, sulfur, etc. are fixed, thereby improving weldability of steel; secondly, due to the effect of microscopic particles, such as the insolubility of V (CN) at high temperature, coarsening of the crystal grains in the heat affected zone is prevented, and the toughness of the heat affected zone is improved, thereby improving the weldability of the steel. The microalloy elements are beneficial and unnecessary addition elements, and the addition amount is not excessive in consideration of factors such as cost increase and the like.

The invention finely controls the recovery, recrystallization and phase change processes of the deformed structure of the rolled hard strip steel in the heat treatment process by a rapid heat treatment method (comprising the processes of rapid heating, short-time heat preservation and rapid cooling), and finally obtains various fine, uniform and dispersedly distributed tissue structures and good strong plasticity matching.

The specific principle is as follows: different heating rates are adopted in different temperature stages of the heating process, the low-temperature stage mainly recovers deformed tissues, and a relatively low heating rate can be adopted to reduce energy consumption; in the high temperature section, recrystallization and grain growth of different phase structures mainly occur, and relatively high heating rate is needed to shorten the retention time of the structures in the high temperature section so as to ensure grain refinement. The recovery of a deformed structure and a ferrite recrystallization process in the heating process are inhibited by controlling the heating rate in the heating process, so that the recrystallization process is overlapped with the austenite phase transformation process, the nucleation points of recrystallized grains and austenite grains are increased, and the grains are refined finally. By short-time heat preservation and quick cooling, the time for grain growth in the soaking process is shortened, and the fine and uniform distribution of grain structures is ensured.

In the heat treatment process disclosed in chinese patent CN107794357B and US patent US2019/0153558A1, although the heating process is also segmented: heating to 300-500 deg.C at a heating rate of 1-10 deg.C/s, heating to 850-950 deg.C at a heating rate of 100-500 deg.C/s, maintaining for no more than 5s, and water-quenching to room temperature. The treatment method requires that the steel plate must be heated to a high-temperature area of single-phase austenite, which improves the high-temperature resistance requirement of equipment and increases the manufacturing difficulty, and simultaneously, a water-cooling mode is adopted, so that although the cooling speed is extremely high, the growth time of a grain structure in the high-temperature area can be greatly reduced, the distribution of alloy elements in a final product is inevitably uneven, the structure performance of the product is uneven and unstable, and the over-high water quenching cooling speed can also cause a series of problems of poor plate shape of the steel plate, surface oxidation and the like.

Only by integrated control of the entire heat treatment process: the method comprises the processes of rapid heating (heating speed is controlled by sections), short-time soaking and rapid cooling, so that the optimal grain size, alloy elements and phase structure which are finely controlled can be obtained and the optimal toughness matching product can be obtained finally.

The main phase structure of the Q & P steel obtained by the rapid heat treatment method of the invention is martensite (80-90% by volume fraction) and residual austenite (10-20% by volume fraction) and contains a very small amount of ferrite (3-5% by volume fraction), so the phase structure is strictly a multi-phase structure, the matrix structure is uniformly distributed, obvious lamellar tempered martensite appears, the grain diameter is 1-3 μm, uniformly distributed ferrite phase exists around the martensite strengthening phase grains, and the martensite strengthening phase grains are mainly in a lamellar structure.

The invention relates to a rapid heat treatment manufacturing method of 1280 MPa-grade low-carbon low-alloy Q & P steel, which comprises the following steps:

1) Smelting and casting

Smelting according to the chemical components and casting into a plate blank;

2) Hot rolling and coiling

The coiling temperature is 550-680 ℃;

3) Cold rolling of steel

Cold rolling reduction rate is 40-85%, and rolling hard state strip steel or steel plate is obtained;

4) Rapid thermal processing

a) Rapid heating

Rapidly heating the cold-rolled strip steel or the steel plate from room temperature to a target temperature of an austenite and ferrite two-phase region at 770-845 ℃; the rapid heating adopts a one-stage type or two-stage type, the heating rate is 50-500 ℃/s when the one-stage type rapid heating is adopted, when the two-stage type rapid heating is adopted, the first stage is heated to 550-625 ℃ from room temperature at the heating rate of 15-500 ℃/s, and the second stage is heated to 770-845 ℃ from 550-625 ℃ at the heating rate of 50-500 ℃/s;

b) Soaking heat

Soaking at the target temperature of 770-845 ℃ in an austenite and ferrite two-phase region for 10-60 s;

c) Cooling down

After the heat equalization of the strip steel or the steel plate is finished, slowly cooling the strip steel or the steel plate to 700-770 ℃ at a cooling rate of 5-15 ℃/s, then quickly cooling the strip steel or the steel plate to 230-280 ℃ at a cooling rate of 50-200 ℃/s, and preserving heat for 2-10 s in the temperature interval;

d) Tempering

After the heat preservation is finished, the strip steel or the steel plate is heated to 300-470 ℃ at the heating rate of 10-30 ℃/s for tempering treatment, and the tempering time is 10-60 s;

e) And after the tempering is finished, cooling the strip steel or the steel plate to the room temperature at the cooling rate of 30-100 ℃/s.

Preferably, in the step 2), the hot rolling temperature is not less than Ar ₃ 。

Preferably, in the step 2), the coiling temperature is 580 to 650 ℃.

Preferably, in the step 3), the cold rolling reduction is 60 to 80%.

Preferably, the rapid heating is performed in a one-stage heating mode, and the heating rate is 50-300 ℃/s.

Preferably, the rapid heating adopts two-stage heating, and the first stage is heated from room temperature to 550-625 ℃ at the heating rate of 15-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 50-300 ℃/s.

Preferably, the rapid heating adopts two-stage heating, and the first stage is heated from room temperature to 550-625 ℃ at the heating rate of 30-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 80-300 ℃/s.

Preferably, the rapid heating final temperature is 790 to 845 ℃.

Preferably, in the soaking process in the step 4), after the strip steel or the steel plate is heated to the target temperature of the two-phase region of austenite and ferrite, soaking is carried out while keeping the temperature unchanged, and the soaking time is 10-40 s.

Preferably, in the soaking process in the step 4), the temperature of the strip steel or the steel plate is raised or lowered within a small range within the soaking time period, the temperature after raising is not more than 845 ℃, the temperature after lowering is not less than 770 ℃, and the soaking time is 10-40 s.

1. heating rate control

Generally, under the traditional slow heating condition, a deformation matrix firstly recovers, recrystallizes and grows grains, then phase transformation from ferrite to austenite occurs, phase transformation nucleation is mainly performed at the grain boundary of the grown ferrite, the nucleation rate is low, and finally the obtained material texture grains are relatively coarse. The recrystallization kinetics of the continuous heating process can be quantitatively described by the relationship affected by the heating rate, the volume fraction of ferrite recrystallized during continuous heating as a function of temperature T:

wherein X (t) is ferrite recrystallization volume fraction; n is an Avrami index, is related to a phase change mechanism, depends on the decay period of the recrystallization nucleation rate, and generally takes a value within the range of 1-4; t is the heat treatment temperature; t is _star Is the recrystallization onset temperature; β is the heating rate; b (T) is obtained by the following formula:

b＝b ₀ exp(-Q/RT)

it can be derived from the above formula and the related experimental data that the recrystallization onset temperature (T) increases with the rate of heating _star ) And end temperature (T) _fin ) All rise; when the heating rate is more than 50 ℃/s, the austenite phase transformation and the recrystallization process are overlapped, the recrystallization temperature is raised to the temperature of the two-phase region, and the faster the heating rate is, the ferriteThe higher the recrystallization temperature.

Under the rapid heating condition, the deformed matrix does not complete recrystallization (even does not fully recover) or just completes recrystallization, the phase transformation from ferrite to austenite begins to occur, and because the grains are fine and the grain boundary area is large when the recrystallization is just completed, the nucleation rate is obviously improved, and the grains are obviously refined. Particularly, after the ferrite recrystallization process and the austenite phase transformation process are overlapped, a large number of crystal defects such as dislocation and the like are reserved in the deformed ferrite crystal at the part which is not completely recrystallized, a large number of nucleation points are provided for austenite, and the nucleation of the austenite presents an explosive nucleation mode, so that the austenite crystal grains are further refined, and the high-density dislocation line defects also become channels for high-speed diffusion of carbon atoms, so that each austenite crystal grain can be quickly generated and grown, and the volume fraction of the austenite is increased. The rapid heating process lays a good foundation for the transformation from austenite to martensite in the subsequent rapid cooling process.

The invention comprehensively considers the factors of the effect of rapidly heating and thinning crystal grains, the manufacturing cost, the manufacturability and the like, and sets the heating rate to be 50-500 ℃/s when one-stage rapid heating is adopted and 15-500 ℃/s when two-stage rapid heating is adopted.

Because the influence of rapid heating on the structure evolution processes of material such as recovery, recrystallization and grain growth is different in different temperature interval ranges, in order to obtain optimal structure control, the optimal heating rates of different heating temperature intervals are different: the heating rate has the greatest influence on the recovery process from 20 ℃ to 500-625 ℃, and is controlled to be 15-300 ℃/s, and is further preferably 30-300 ℃/s; the heating temperature is 500-625 ℃ to the austenitizing temperature 770-845 ℃, the heating rate has the greatest influence on the processes of recrystallization nucleation, phase change nucleation and grain growth, and the heating rate is controlled to be 50-300 ℃/s; more preferably 80 to 300 ℃/s.

2. Soaking temperature control

The soaking temperature depends on the content of C, and the soaking temperature is generally set at A in the traditional process _C3 Above 30-50 ℃, the hairObviously, a great deal of dislocation is formed in ferrite by utilizing a rapid heating technology to provide nucleation work for austenite transformation, so that the temperature is only required to be heated to A _C1 To A _C3 In between. The C content of the Q & P steel is 0.16-0.23%, A _C1 And A _C3 About 730 ℃ and 870 ℃ respectively. There are a large amount of undissolved tiny evenly distributed carbide in Q & P steel, in the soaking process, can play the effect of mechanical hindrance to the growth of austenite granule, be favorable to refining the grain size of alloy steel, but if heating temperature is too high, will make undissolved carbide figure reduce in a large number, weaken this kind of hindrance effect, strengthen the growing up tendency of grain, and then reduce the intensity of steel. When the amount of undissolved carbides is too large, aggregation may occur, resulting in nonuniform distribution of local chemical components, and when the carbon content in the aggregated portion is too high, local overheating may also occur. Ideally, a small amount of fine granular undissolved carbides should be uniformly distributed in the alloy steel, so that not only can the abnormal growth of austenite grains be prevented, but also the content of each alloy element in a matrix can be correspondingly increased, and the aim of improving the mechanical properties of the alloy steel, such as strength, toughness and the like, is fulfilled.

The soaking temperature is selected with the aim of obtaining fine and uniform austenite grains, so as to achieve the final aim of obtaining a fine and uniform martensite structure with a higher volume fraction after cooling. The overhigh soaking temperature can cause austenite grains to be coarse, workpieces are easy to crack in the quenching process, and the martensite structure obtained after quenching is also coarse, so that the mechanical property of the steel is poor; but also reduce the amount of retained austenite in the Q & P steel, and reduce the hardness and wear resistance of the material. Too low soaking temperature can lead the content of carbon and alloy elements dissolved in austenite to be insufficient, lead the carbon concentration of austenite to be unevenly distributed, greatly reduce the hardenability of steel and cause adverse effect on the mechanical property of steel. The soaking temperature of the hypoeutectoid steel should be Ac ₃ +30 to 50 ℃. In the case of ultra-high strength steels, the presence of carbide-forming elements hinders the transformation of carbides, so the soaking temperature can be suitably increased. Combining the above factors, the invention selects 770-850 ℃ as soaking temperature to obtain more rationalityA more reasonable final organization is desired.

3. Soaking time control

Because the rapid heating is adopted, the material in the two-phase region contains a large amount of dislocation, a large amount of nucleation points are provided for the formation of austenite, and a rapid diffusion channel is provided for carbon atoms, so that the austenite can be formed very rapidly, and the shorter the soaking and heat-preserving time is, the shorter the diffusion distance of the carbon atoms is, the larger the carbon concentration gradient in the austenite is, and the more the carbon content of the residual austenite is remained; however, if the heat preservation time is too short, the distribution of alloy elements in the steel is uneven, and the austenitizing is insufficient; too long heat preservation time easily causes coarse austenite grains. The influence factor of the soaking and heat-preserving time also depends on the contents of carbon and alloy elements in the steel, when the contents of the carbon and the alloy elements in the steel are increased, the thermal conductivity of the steel is reduced, and because the diffusion speed of the alloy elements is slower than that of the carbon elements, the alloy elements obviously delay the structure transformation of the steel, and the heat-preserving time is properly prolonged. In conclusion, the soaking and heat preservation time is set to be 10-60 s.

4. Rapid cooling rate control

The control of the rapid cooling process needs to be combined with comprehensive factors such as the evolution result of each phase structure and the alloy diffusion distribution result in the early heating and soaking processes, and the like, so as to ensure that the ideal material structure with each phase structure and reasonably distributed elements is finally obtained.

In order to obtain a martensite strengthening phase, the cooling speed of the material during rapid cooling is required to be greater than the critical cooling speed so as to obtain a martensite structure, the critical cooling speed mainly depends on the material components, the content of Si in the invention is 1.4-2.0%, the content of Mn is 2.4-3.0%, and the content is relatively high, so that the hardenability of Q & P steel is greatly enhanced by Si and Mn, and the critical cooling speed is reduced. The cooling rate is too low to obtain martensite structure, and the mechanical property can not meet the requirement; too high cooling rate can generate larger quenching stress (namely, structural stress and thermal stress) to cause serious defects of plate shape, and easily cause deformation and cracking of the Q & P steel plate strip. Therefore, the rapid cooling speed is set to be 50-200 ℃/s.

5. Tempering temperature control

Generally, when alloy steel is tempered below 150 ℃, because the temperature is too low, alloying elements cannot diffuse, and only carbon element has certain diffusion capacity, so that the low-temperature tempered steel has high hardness, but has too high brittleness and poor toughness, and cannot meet the service performance requirements of workpieces. When the tempering is performed at a temperature of 200 ℃ or higher, a large amount of carbon elements and other alloying elements contained in martensite start to be precipitated, so that the residual stress is reduced to be eliminated, and the hardness of the tempered steel gradually decreases with the increase of the tempering temperature, but the toughness is enhanced. When the tempering temperature reaches about 500 ℃, the martensite decomposition is finished, cementite gradually aggregates and grows, the alpha phase starts to generate a recovery process, the temperature is continuously raised, the alpha phase starts to recrystallize to form polygonal ferrite, and the strength is obviously reduced. The higher the tempering temperature is, the coarser the alpha phase and the carburized phase are, and the lower the hardness of the tempered steel is, and the final purpose of the invention is to obtain better strength and plasticity at the same time, so the tempering temperature is set to be 300-470 ℃.

6. Tempering time control

In the tempering process of the steel, the tempering time plays three roles: (1) ensuring that the tissue transformation is fully performed; (2) reducing or eliminating internal stress; (3) The workpiece obtains the required performance by matching with the tempering temperature. Because the steel of the invention adopts the rapid heating technology to refine austenite grains, the distance between the retained austenite and the martensite generated after the primary rapid cooling is shortened, the efficiency of the diffusion and distribution of carbon atoms from the supersaturated martensite to the retained austenite is improved, and the time required by the tempering process is greatly reduced. But if the tempering time is too short, the internal stress is difficult to eliminate, and the brittleness and the hardness of the workpiece are reduced, and the tempering time is set to be 10-60 s comprehensively.

The invention realizes the rapid heat treatment process by carrying out rapid heating and rapid cooling process transformation on the traditional continuous annealing unit, can greatly shorten the length of a heating and soaking section of the annealing furnace (at least one third of the length of the traditional continuous annealing furnace), improves the production efficiency of the traditional continuous annealing unit, reduces the production cost and energy consumption, obviously reduces the number of furnace rollers of the continuous annealing furnace, particularly the number of furnace rollers of a high-temperature furnace section, can improve the surface quality control capability of strip steel, and obtains the strip steel product with high surface quality. Meanwhile, by establishing a novel continuous annealing unit adopting a rapid heat treatment process technology, the purposes of short and bold unit, flexible material transition, strong regulation and control capability and the like can be realized; for the product material, the grain of the strip steel can be refined, the strength and the plasticity of the material are further improved, the alloy cost and the manufacturing difficulty of the working procedure before heat treatment are reduced, and the use performance of the material for users such as forming, welding and the like is improved.

Compared with the prior art, the invention has the advantages that:

(1) The invention inhibits the recovery of a deformed structure and a ferrite recrystallization process in a heat treatment process by rapid heat treatment, so that the recrystallization process is overlapped with an austenite phase transformation process, the nucleation points of recrystallized grains and austenite grains are increased, the grain growth time is shortened, the grains are refined, the obtained multiphase structure of which the metallographic structure martensite accounts for 80-90%, the residual austenite accounts for 10-20% and the ferrite accounts for 3-5% of the Q & P steel is uniform in matrix structure distribution, obvious lamellar tempered martensite appears, the grain size is refined to 1-3 mu m, a uniformly distributed ferrite phase exists around the martensite strengthening phase grains, and the martensite strengthening phase grains are mainly of a lamellar structure; the austenite in the structure has various forms such as blocks, strips, particles and the like, has good thermal stability, has austenite transformation rate of below 8 percent at 50 ℃ below zero and austenite transformation rate of below 30 percent at 190 ℃ below zero, and can continuously generate TRIP effect under different strain conditions, so that the mechanical property and the user service performance of the product are excellent.

(2) Compared with Q & P steel obtained by the traditional heat treatment mode, the alloy components of the Q & P steel obtained by the invention are greatly reduced, the grain size is reduced by 40-70%, the yield strength is 754-1112 MPa, the tensile strength is 1281-1350 MPa, the elongation is 19-22.2%, and the product of strength and elongation is 24.8-28.97 GPa%.

(3) According to the low-carbon low-alloy Q & P steel rapid heat treatment process of 1280MPa grade, the time for the whole heat treatment process can be shortened to 71-186 s, the time of the whole heat treatment process is greatly reduced (the time of the traditional continuous annealing process of the Q & P steel is usually 5-8 min), the production efficiency is improved, the energy consumption is reduced, and the production cost is reduced.

(4) Compared with the traditional Q & P steel and the heat treatment process thereof, the rapid heat treatment method shortens the time of the heating section and the soaking section by 60-80%, shortens the time of the whole heat treatment process to 71-186 s, can save energy, reduce emission and consumption, remarkably reduces the one-time investment of equipment such as a furnace and the like, and remarkably reduces the production running cost and the equipment maintenance cost; in addition, the alloy content can be reduced by producing products with the same strength grade through rapid heat treatment, the production cost of heat treatment and previous processes is reduced, and the manufacturing difficulty of each process before heat treatment is reduced.

(5) Compared with the Q & P steel produced by the traditional process and the heat treatment process thereof, the rapid heat treatment process technology shortens the time of the heating process and the soaking process, shortens the length of the furnace, reduces the number of furnace rollers and reduces the probability of surface defects of Q & P steel strip steel products in the furnace, so the surface quality of the products is obviously improved; in addition, due to the refinement of product crystal grains and the reduction of material alloy content, the processing forming performance such as hole expanding performance, bending performance and the like, and the user service performance such as welding performance and the like of the Q & P steel obtained by adopting the technology of the invention are also improved.

The low-carbon low-alloy 1280 MPa-grade Q & P steel obtained by the invention has important value for the development of new-generation light-weight transportation tools such as automobiles, trains, ships, airplanes and the like and the healthy development of corresponding industries and advanced manufacturing industries.

Drawings

FIG. 1 is a photograph of the microstructure of Q & P steel produced in example 1 of test steel A of the present invention.

FIG. 2 is a photograph showing the microstructure of Q & P steel produced by conventional process 1 of test steel A of the present invention.

FIG. 3 is a photograph of the microstructure of Q & P steel produced in example 7 of test steel K of the present invention.

FIG. 4 is a photograph of the microstructure of Q & P steel produced in example 8 of test steel R of the present invention.

FIG. 5 is a photograph showing the microstructure of Q & P steel produced in example 22 of test steel P of the present invention.

FIG. 6 is a photograph showing the microstructure of Q & P steel produced in example 23 of test steel S of the present invention.

Detailed Description

The present invention is further illustrated with reference to the following examples and the accompanying drawings, wherein the examples are carried out on the premise of the technical solution of the present invention, and detailed embodiments and specific operating procedures are given, but the scope of the present invention is not limited to the following examples.

The compositions of the test steels according to the present invention are shown in Table 1, the specific parameters of the examples according to the present invention and the conventional process are shown in tables 2 and 3, and tables 4 and 5 show the main properties of the steels prepared according to the examples and the conventional process.

As can be seen from tables 1 to 4, by the method of the present invention, the alloy content in the steel of the same grade can be reduced, the crystal grains are refined, and the good matching of the material structure composition and the strength and the toughness is obtained. The Q & P steel obtained by the method has the yield strength of 754-1112 MPa, the tensile strength of 1281-1350 MPa, the elongation of 19-22.2 percent and the product of strength and elongation of 24.8-28.97 GPa percent.

FIG. 1 is a structural diagram of a typical composition A steel obtained in example 1, and FIG. 2 is a structural diagram of a typical composition A steel obtained in conventional process example 1. From the figure, the material structures treated by different heat treatment modes are very different. The structure of the steel obtained by the processing of the embodiment of the invention mainly comprises fine and uniform martensite structures and a small amount of carbides which are dispersed and distributed on a ferrite matrix, and the martensite grain structures and the small amount of carbides are very fine and uniformly distributed in the ferrite matrix, which is very favorable for improving the strength and the plasticity of the material. The steel structure treated by the traditional process is relatively non-uniform in distribution, the martensite crystal grains are relatively large, and a small amount of residual austenite and carbide structures are distributed on the martensite crystal boundary and are non-uniform in distribution. The tissue characteristics treated by the traditional process are as follows: the crystal grains are relatively coarse and have a certain uneven distribution of the structure.

FIG. 3 is a structural diagram of a typical composition K steel obtained in example 7, and FIG. 4 is a structural diagram of a typical composition R steel obtained in example 8. FIG. 5 is a structural diagram of a typical composition P steel obtained in example 22, and FIG. 6 is a structural diagram of a typical composition S steel obtained in example 23. Examples 7, 8, 22 and 23 are all processes with a short overall heat treatment period. As can be seen from the figure, by adopting the method of the invention, more uniform, fine and dispersedly distributed phase structures can be obtained through short-time rapid annealing treatment. Therefore, the preparation method of the invention can refine the crystal grains, and lead each phase structure of the material to be evenly distributed in the matrix, thereby improving the material structure and the material performance.

The invention can transform the traditional continuous annealing unit by adopting the rapid heating and rapid cooling process, so that the rapid heat treatment process is realized, the length of the heating and soaking sections of the traditional continuous annealing furnace can be greatly shortened, the production efficiency of the traditional continuous annealing unit is improved, the production cost and the energy consumption are reduced, the number of furnace rollers of the continuous annealing furnace is reduced, the control capability of the surface quality of the strip steel can be improved, and the strip steel product with high surface quality is obtained; meanwhile, by establishing a novel continuous annealing unit adopting a rapid heat treatment process technology, the continuous heat treatment unit has the advantages of short and concise structure, flexible material transition, strong regulation and control capability and the like; for the material, the grain of the strip steel can be refined, the strength of the material is further improved, the alloy cost and the manufacturing difficulty of the working procedure before heat treatment are reduced, and the welding performance and other user service performances of the material are improved.

In conclusion, the invention adopts the rapid heat treatment process, so that the technological progress of the continuous annealing process of the cold-rolled strip steel is greatly promoted, the austenitizing process of the cold-rolled strip steel from room temperature to the last is expected to be completed within dozens of seconds, dozens of seconds or even several seconds, the heating section length of the continuous annealing furnace is greatly shortened, the speed and the production efficiency of a continuous annealing unit are convenient to improve, the number of rollers in the furnace of the continuous annealing unit is obviously reduced, for a rapid heat treatment production line with the unit speed of about 180 meters/minute, the number of rollers in the high-temperature furnace section is not more than 10, and the surface quality of the strip steel can be obviously improved. Meanwhile, the rapid heat treatment process of the recrystallization and austenitizing processes finished in a very short time also provides a more flexible and flexible high-strength steel structure design method, so that the material structure is improved and the material performance is improved on the premise of not changing alloy components, rolling process and other pre-process conditions.

The advanced high-strength steel represented by Q & P steel has wide application prospect, the rapid heat treatment technology has great development and application value, and the combination of the two technologies can provide larger space for the development and production of the Q & P steel.

Claims

1. A1280 MPa-grade low-carbon low-alloy Q & P steel comprises the following chemical components in percentage by mass: c:0.16 to 0.23%, si:1.4 to 2.0%, mn:2.4 to 3.0%, ti: 0.006-0.016%, P is less than or equal to 0.015%, S is less than or equal to 0.002%, al: 0.02-0.05%, and one or two of Cr, mo, nb and V, wherein Cr + Mo + Ti + Nb + V is less than or equal to 0.5%, and the balance of Fe and other unavoidable impurities, and is obtained by the following process:

1) Smelting and casting

Smelting according to the chemical components and casting into a plate blank;

2) Hot rolling and coiling

The coiling temperature is 550-680 ℃;

3) Cold rolling

The cold rolling reduction rate is 40-85%;

4) Rapid thermal processing

Rapidly heating the cold-rolled steel plate to 770-845 ℃, wherein the rapid heating adopts a one-stage or two-stage mode, and the heating rate is 50-500 ℃/s when the one-stage rapid heating is adopted; when two-stage rapid heating is adopted, the first stage is heated from room temperature to 550-625 ℃ at the heating rate of 15-500 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at the heating rate of 50-500 ℃/s; then soaking for 10-60 s at 770-845 ℃;

slowly cooling to 700-770 ℃ at a cooling rate of 5-15 ℃/s after heat equalization, then rapidly cooling to 230-280 ℃ at a cooling rate of 50-200 ℃/s, preserving heat for 2-10 s in the temperature interval, then heating to 300-470 ℃ at a heating rate of 10-30 ℃/s for tempering, and tempering for 10-60 s; after the tempering is finished, cooling to room temperature at a cooling rate of 30-100 ℃/s.

2. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein the C content is 0.18-0.21%.

3. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein the Si content is 1.6-1.8%.

4. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein the Mn content is 2.6-2.8%.

5. The 1280MPa grade low-carbon low-alloy Q of claim 1&P steel, characterized in that in the step 2), the hot rolling temperature is more than or equal to A _r3 。

6. The 1280MPa grade low carbon low alloy Q & P steel as claimed in claim 1 or 5, wherein the coiling temperature in step 2) is 580-650 ℃.

7. The low-carbon low-alloy Q & P steel of 1280MPa grade according to claim 1, wherein the cold rolling reduction in step 3) is 60-80%.

8. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein the time for the entire rapid heat treatment process is 71-186 s.

9. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein the rapid heating in step 1) is performed in one stage at a heating rate of 50-300 ℃/s.

10. The 1280MPa grade low carbon low alloy Q & P steel according to claim 1, wherein the rapid heating of step 1) is performed in two stages: the first section is heated from room temperature to 550-625 ℃ at the heating rate of 15-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 50-300 ℃/s.

11. The 1280MPa grade low carbon low alloy Q & P steel according to claim 1, wherein the rapid heating of step 1) is performed in two stages: the first section is heated from room temperature to 550-625 ℃ at the heating rate of 30-300 ℃/s; the second section is heated from 550-625 ℃ to 790-845 ℃ at the heating rate of 80-300 ℃/s.

12. The 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 1, wherein in the step 4), the strip steel or the steel plate is rapidly cooled from 700-770 ℃ to 230-280 ℃ at a cooling rate of 50-150 ℃/s.

13. The 1280 MPa-level low-carbon low-alloy Q & P steel according to any of claims 1 to 12, wherein the Q & P steel has a metallographic structure comprising 80 to 90% of martensite, 10 to 20% of retained austenite and 3 to 5% of ferrite, and has a uniform matrix structure, a significantly lamellar tempered martensite appears, a grain size of 1 to 3 μm, a uniformly distributed ferrite phase exists around martensite reinforcing phase grains, and the martensite reinforcing phase grains are mainly in a lamellar structure.

14. The 1280 MPa-grade low-carbon low-alloy Q & P steel according to any of claims 1-13, wherein said Q & P steel has good thermal stability of austenite in metallographic structure, austenite transformation rate at-50 ℃ is lower than 8%, austenite transformation rate at-190 ℃ is lower than 30%.

15. The low-carbon low-alloy Q & P steel with the grade of 1280MPa according to any one of claims 1 to 14, wherein the Q & P steel has the yield strength of 754 to 1112MPa, the tensile strength of 1281 to 1350MPa, the elongation of 19 to 22.2 percent and the product of strength and elongation of 24.8 to 28.97GPa%.

16. The rapid thermal processing manufacturing method of 1280MPa grade low carbon low alloy Q & P steel according to any of claims 1 to 15, characterized by comprising the following steps:

1) Smelting and casting

Smelting according to the chemical components and casting into a plate blank;

2) Hot rolling and coiling

The coiling temperature is 550-680 ℃;

3) Cold rolling of steel

The cold rolling reduction rate is 40-85%, and the rolling hard strip steel or steel plate is obtained after cold rolling;

4) Rapid thermal processing

a) Rapid heating

Rapidly heating the cold-rolled strip steel or the steel plate from room temperature to a target temperature of an austenite and ferrite two-phase region at 770-845 ℃; the rapid heating adopts a one-stage or two-stage type, the heating rate is 50-500 ℃/s when the one-stage type rapid heating is adopted, when the two-stage type rapid heating is adopted, the first stage is heated from room temperature to 550-625 ℃ at the heating rate of 15-500 ℃/s, and the second stage is heated from 550-625 ℃ to 770-845 ℃ at the heating rate of 50-500 ℃/s;

b) Soaking heat

c) Cooling down

d) Tempering

After the heat preservation is finished, heating the strip steel or the steel plate to 300-470 ℃ at the heating rate of 10-30 ℃/s for tempering for 10-60 s;

17. The 1280MPa grade low-carbon low-alloy Q of claim 16&The rapid heat treatment manufacturing method of the P steel is characterized in that in the step 2), the hot rolling temperature is more than or equal to Ar ₃ 。

18. The rapid thermal processing method of manufacturing low-carbon low-alloy Q & P steel of 1280MPa grade according to claim 16 or 17, wherein the coiling temperature in the step 2) is 580 to 650 ℃.

19. The rapid heat treatment manufacturing method of low carbon low alloy Q & P steel of 1280MPa grade according to claim 16, wherein the cold rolling reduction in step 3) is 60-80%.

20. The rapid heat treatment manufacturing method of 1280MPa grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein the time for the whole process of the rapid heat treatment is 71-186 s.

21. The rapid thermal processing method of 1280MPa grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein the rapid heating is performed in one stage at a heating rate of 50-300 ℃/s.

22. The rapid thermal processing manufacturing method of 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein said rapid heating is performed in two stages, the first stage is heated from room temperature to 550-625 ℃ at a heating rate of 15-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 50-300 ℃/s.

23. The rapid thermal processing manufacturing method of 1280MPa grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein the rapid heating adopts two-stage heating, the first stage is heated from room temperature to 550-625 ℃ at a heating rate of 30-300 ℃/s; the second section is heated from 550-625 ℃ to 770-845 ℃ at a heating rate of 80-300 ℃/s.

24. The rapid thermal processing manufacturing method of 1280MPa grade low carbon low alloy Q & P steel of claim 16, wherein the rapid heating final temperature is 790 to 845 ℃.

25. The rapid thermal processing method of 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein the soaking time is 10-40 s, after the strip steel or the steel plate is heated to the target temperature of the two-phase region of austenite and ferrite in the soaking process of step 4).

26. The rapid thermal processing method of 1280 MPa-grade low-carbon low-alloy Q & P steel as claimed in claim 16, wherein in the soaking process of step 4), the strip steel or the steel plate is heated up or cooled down with small amplitude within the soaking time period, the temperature after heating up is not more than 845 ℃, the temperature after cooling down is not less than 790 ℃, and the soaking time is 10-40 s.

27. The rapid thermal processing method of manufacturing low-carbon low-alloy Q & P steel of 1280MPa grade according to claim 16, wherein the strip or the plate is rapidly cooled from 700 to 770 ℃ to 230 to 280 ℃ at a cooling rate of 50 to 150 ℃/s in step 4).