A process for improving the cold stamping formability of 1000mpa grade automotive steel

Abstract

The application discloses a process for improving the cold stamping forming performance of 1000MPa grade automobile steel, which comprises the following steps: placing medium-manganese steel raw materials in a vacuum induction furnace for smelting, adding rare earth iron alloy after the medium-manganese steel raw materials are completely melted, uniformly treating after the rare earth iron alloy is completely melted, then smelting and casting to obtain a blank; heating and keeping the blank, then hot rolling to obtain a hot-rolled plate; soaking the hot-rolled plate in hydrochloric acid, then washing the hydrochloric acid on the surface of the hot-rolled plate with clean water; then blowing dry and coating oil on the surface of the hot-rolled plate to obtain a pickled hot-rolled plate; annealing and softening the pickled hot-rolled plate, then cold rolling to obtain a cold-rolled plate; and annealing the cold-rolled plate. In the application, rare earth elements are added to the medium-manganese steel, and the best heat treatment process and deformation process parameters are determined through detection and experiments, so that the super-high-strength automobile steel with better strength and plasticity is obtained, and the production cost is effectively controlled.

Description

TECHNICAL FIELD

[0001] The present application relates to the technical field of 1000MPa grade or above ultra-high strength automobile steel. Specifically, it is a process for improving the cold stamping forming performance of 1000MPa grade automobile steel. BACKGROUND

[0002] The third generation automobile steel is a medium manganese steel mainly with BCC structure ferrite and a small amount of FCC structure austenite, which has excellent mechanical properties and forming performance, and is widely used in the field of automobile manufacturing. The good characteristics of the microstructure, mechanical properties and forming performance of the third generation automobile steel are due to the presence of a suitable amount of metastable residual austenite at room temperature.

[0003] Ultra-high strength automobile steel needs to meet the requirements of low cost, high plasticity and high strength, and the key to improving ultra-high strength automobile steel is to improve the heat treatment process, composition design and rolling process.

[0004] At present, the mainstream heat treatment process of the third generation automobile steel includes austenite reverse transformation (ART) process and two-phase zone direct annealing (IA) process. The process, heat treatment temperature, holding time, heating rate and cooling rate of these two processes will affect the microstructure and properties of the steel.

[0005] CN113549745B discloses a third generation automobile steel processing process. In this process, for a test steel with a C content of 0.12wt%, a Mn content of 5.2wt%, a Si content of 0.04wt%, an Al content of 0.04wt%, a Cu content of 0.24wt%, a Ni content of 0.25wt%, a Nb content of 0.025wt%, a Ti content of 0.022wt%, a P content of 0.011wt%, a S content of 0.004wt%, and a N content of 0.005wt%, the remainder being Fe, the test steel is first hot-rolled, then cold-rolled, and an independent intermediate annealing process is added during cold rolling. The intermediate annealing conditions are: annealing temperature 500-580℃, annealing time 20-40min, and cold rolling reduction 55%-80%. After cold rolling, two-phase zone direct annealing or ART annealing process is adopted. The ART annealing method is: first quenching at 780-820℃ for 5-10min, then holding at 620-640℃ for 5-15min, and then cooling to room temperature by air cooling at a speed of 1-2℃ / s. The two-phase zone direct annealing method is: holding at 620-640℃ for 5-15min, and then cooling to room temperature by air cooling at a speed of 1-2℃ / s. The elongation of the experimental steel prepared by two-phase zone direct annealing process is lower than that of the experimental steel prepared by ART annealing process, but the tensile strength, yield strength and strength plasticity product of the experimental steel prepared by two-phase zone direct annealing process are higher than those of the experimental steel prepared by ART annealing process, and the experimental steel prepared by two-phase zone direct annealing process has better deep drawing performance.

[0006] The deformation process of steel also affects its microstructure and properties. In the paper "Effect of Cold Rolling Reduction on Microstructure and Forming Properties of Low-cost Third Generation Automobile Steel", Dong Ruifeng et al. studied the effect of cold rolling reduction on the microstructure and forming properties of third generation automobile steel. The chemical composition of the experimental steel is shown in Table 1, and the experimental results are shown in Tables 2 and 3.

[0007] Table 1 Chemical composition of experimental steel (%)

[0008]

[0009] Table 2 Volume fraction of texture of experimental steel under different cold rolling reductions

[0010]

[0011] Table 3 Mechanical properties of experimental steel under different cold rolling reductions after heat treatment

[0012]

[0013] The experimental results show that after annealing at 630℃ for 10min, the microstructure of experimental steel with cold rolling reductions of 66%, 73% and 79% is similar, and the conventional mechanical properties are similar. The microstructure is mainly austenite and ferrite. With the increase of reduction, the texture components of {111} and {110} that are beneficial to stamping forming performance increase, while the {100} texture component decreases, and the plastic strain ratio r increases with the increase of reduction. When the reduction is 79%, the content of {111} and {110} favorable texture is the highest, which is 17.7% and 25.5% respectively, and the r value is the largest, which is 0.961. The experimental steel has good forming performance.

[0014] The content of rare earth elements in steel affects the texture of steel. In the paper "Research on Annealing Process of Rare Earth Micro-alloyed Low-cost Third Generation Automobile Steel", Zhao Qingbo et al. compared the texture content of two medium manganese steels with and without rare earth under two-phase region direct annealing and ART annealing process. The results are shown in the table. Table 4 shows the chemical composition of the two experimental steels used for comparison, and Tables 5 and 6 show the texture content of each orientation under two-phase region direct annealing and ART annealing process respectively.

[0015] Table 4 Chemical composition of experimental steel (%)

[0016]

[0017]

[0018] Table 5 Texture content of each orientation under ART annealing process

[0019]

[0020] Table 6 Texture content of each orientation under two-phase zone direct annealing process

[0021]

[0022] The texture content of {111}, {110} and {100} planes of two groups of experimental steels under two kinds of heat treatment processes is calculated and analyzed by using ResMat-TexTools software, as shown in the following table. Figure 1 As can be seen from the figure, for the ART annealing process, the {110} plane texture content of the experimental steel B with trace rare earth elements is 19.6%, the {111} plane texture content is 13.4%, and the content of the unfavorable texture {100} plane is 11.1%; for the two-phase zone direct annealing process, the {110} plane texture content of the experimental steel B containing rare earth Ce is 23.5%, the {111} plane texture content is 16.9%, and the content of the {100} plane texture is 10.0%. By comparison, it is found that no matter which heat treatment process is used, the {111} and {110} texture contents of the experimental steel B containing trace rare earth element Ce are slightly higher than those of the experimental steel A without rare earth elements, and the content of the {100} texture is slightly lower.

[0023] In addition to the texture content, the forming performance of the steel is also related to the product of strength and plasticity and the r value. Specifically, for cold stamping forming performance, the steel with a larger product of strength and plasticity usually has better plastic deformation capacity, can withstand a larger deformation amount, and is not prone to cracking, wrinkling and other phenomena; the r value is an important parameter for evaluating the deep drawing performance of a metal sheet, which reflects the ability of a metal sheet to resist thinning or thickening when subjected to tension or pressure in a certain plane, and the better the cold stamping forming performance of the steel, the higher the r value. For the existing steel containing 9ppm of rare earth elements, the highest product of strength and plasticity is 32.90GPa·%, and the cold stamping forming performance still needs to be further improved. SUMMARY

[0024] Therefore, the technical problem to be solved by the present application is to provide a process for improving the cold stamping forming performance of 1000MPa grade automobile steel, by adding rare earth elements to medium manganese steel and controlling the heat treatment process and deformation process parameter conditions, so that the obtained automobile steel has excellent cold stamping forming performance, while effectively controlling the production cost.

[0025] To solve the above technical problems, the present application provides the following technical solutions:

[0026] A processing process for improving the cold stamping forming performance of automobile steel, comprising the following steps:

[0027] Step A: Melting and Casting: Medium-manganese steel raw material is placed in a vacuum induction furnace for melting. After the medium-manganese steel raw material is completely melted, rare earth ferroalloy is added. After the rare earth ferroalloy is completely melted, homogenization treatment is performed, followed by melting and casting to obtain a billet. Since the amount of alloying elements added is relatively small, if they are melted together with the medium-manganese steel raw material, the rare earth elements in the billet may be completely oxidized and float to the surface of the molten steel as inclusions, failing to remain in the steel to play an alloying role. Therefore, in this invention, cerium ferroalloy is added to the molten steel and homogenized, and thoroughly stirred to ensure that the alloying elements do not segregate, thus ensuring that the rare earth elements are evenly distributed in the billet.

[0028] Step B: Hot rolling: The billet is heated and held at a certain temperature before being hot rolled to obtain a hot-rolled plate;

[0029] Step C: Pickling: After immersing the hot-rolled plate in hydrochloric acid, rinse the surface of the hot-rolled plate with water; then blow dry and apply oil to the surface of the hot-rolled plate to obtain pickled hot-rolled plate; the purpose of pickling is to remove oxide scale, which can cause defects such as pitting on the surface of the steel plate when it is rolled into the steel plate.

[0030] Step D: Cold rolling: First, the pickled hot-rolled sheet is annealed and softened, and then cold-rolled to obtain a cold-rolled sheet;

[0031] Step E: Annealing: The cold-rolled sheet is annealed. After the annealing process, a 1000MPa grade automotive steel with excellent cold stamping performance is obtained.

[0032] The technical solution of the present invention achieves the following beneficial technical effects:

[0033] 1. The present invention optimizes the cold stamping performance of the experimental steel by improving the deformation degree and heat treatment process of medium manganese steel containing rare earth elements. In the hot rolling stage, the experimental steel was subjected to at least 10 hot rolling passes, which improved the degree of hot deformation, made the grains finer, facilitated grain boundary movement, and enhanced the plasticity of the steel, making it easier for the steel to be fully deformed in the cold rolling stage. In the cold rolling stage, by increasing the degree of cold rolling deformation, the grains of the experimental steel were fully elongated and flattened, the dislocation density increased, and the deformation energy of the steel was improved, storing sufficient energy for subsequent heat treatment. In the heat treatment stage, the characteristics of deformation-induced phase transformation and phase transformation-induced plasticity of metastable austenite were utilized to improve the strength-ductility product of the steel. Furthermore, under the influence of rare earth elements on raising the martensitic transformation point and promoting the austenitic transformation, by controlling the cooling rate and holding time, an ideal combination of fine grains and three microstructures—ferrite, martensite, and retained austenite—was obtained, thereby improving both the strength and plasticity of the steel and its cold stamping performance.

[0034] 2. As is well known, increasing the aluminum content is beneficial to improving the ductility and toughness of steel, but the negative effects of increasing the aluminum content must be tolerated. Compared with the prior art disclosed by the inventors' research group of this application (Zhao Qingbo et al., "Research on the Annealing Process of Low-Cost Third-Generation Automotive Steel with Rare Earth Microalloying"), the technical solution of this invention, while reducing the aluminum content by 70%, significantly increases the content of elements such as carbon, manganese, silicon, and cerium, and reduces the niobium content. Furthermore, it increases the number of hot rolling passes to more than 10 in the hot rolling process and controls the annealing temperature in the two-phase region to 640℃. This not only increases the strength-ductility product by up to 22%, but also increases the {111} texture content by 7.3%.

[0035] Compared with the prior art disclosed by the inventors' research group (Dong Ruifeng et al., "The Influence of Cold Rolling Reduction Rate on the Microstructure and Forming Properties of Low-Cost Third-Generation Automotive Steel"), this invention, by adding the rare earth element cerium, increasing the silicon content by more than 6 times, and increasing the content of elements such as manganese and nickel, increases the number of hot rolling passes by more than 2 times. At the same time, it increases the hot rolling start temperature and decreases the final rolling temperature, and increases the annealing temperature and extends the annealing time. Ultimately, this results in an 11% increase in the strength-ductility product of the experimental steel, a 7.3% increase in the {111} texture content, and a 26.8% increase in the plastic strain ratio r value. The cold stamping forming performance of the steel in this invention is significantly improved.

[0036] 3. The processing technology provided by this invention improves the cold forming performance of ultra-high strength automotive steel with a strength of 1000MPa or higher. Under conditions of a bending angle of 0° and a mandrel diameter of 16mm, the steel processed using the technology provided by this invention did not exhibit cracks, indicating that the technology can enable cold-formed steel to achieve the strength of general hot-formed steel, and that it also exhibits excellent cold forming performance. Attached Figure Description

[0037] Figure 1 The texture content of experimental steel under ART annealing and two-phase direct annealing processes in the background art;

[0038] Figure 2a In this embodiment of the invention, JMatPro simulates the CCT results of 0RE experimental steel;

[0039] Figure 2b In this embodiment of the invention, JMatPro simulates the CCT results of 48ppmRE experimental steel;

[0040] Figure 3a The ART annealing process temperature profile used in the embodiments of this invention;

[0041] Figure 3b Temperature profiles of the two-phase direct annealing process used in this embodiment of the invention;

[0042] Figure 4A tensile sample used in the tensile experiment in the embodiment of the present application is shown in the schematic diagram;

[0043] Figure 5a An original austenite morphology diagram of the 0RE experimental steel in the embodiment of the present application;

[0044] Figure 5b An original austenite morphology diagram of the 48ppmRE experimental steel in the embodiment of the present application;

[0045] Figure 5c A grain size distribution diagram of the 0RE experimental steel in the embodiment of the present application;

[0046] Figure 5d A grain size distribution diagram of the 48ppmRE experimental steel in the embodiment of the present application;

[0047] Figure 6a A microstructure diagram of the hot-rolled plate without rare earth elements in the embodiment of the present application;

[0048] Figure 6b A microstructure diagram of the hot-rolled plate containing 48ppm rare earth elements in the embodiment of the present application;

[0049] Figure 7 A test result diagram of the influence of rare earth elements on the conventional mechanical properties of the hot-rolled plate in the embodiment of the present application;

[0050] Figure 8a A microstructure diagram of the cold-rolled plate without rare earth elements in the embodiment of the present application;

[0051] Figure 8b A microstructure diagram of the cold-rolled plate containing 48ppm rare earth elements in the embodiment of the present application;

[0052] Figure 9 A test result diagram of the influence of rare earth elements on the conventional mechanical properties of the cold-rolled plate in the embodiment of the present application;

[0053] Figure 10a An SEM structure diagram of the cold-rolled plate without rare earth elements after quenching in the embodiment of the present application;

[0054] Figure 10b An SEM structure diagram of the hot-rolled plate containing 48ppm rare earth elements after quenching in the embodiment of the present application;

[0055] Figure 11a A microstructure diagram of the 0RE experimental steel under 620℃ ART annealing for 15min in the embodiment of the present application;

[0056] Figure 11b A microstructure diagram of the 0RE experimental steel under 640℃ ART annealing for 15min in the embodiment of the present application;

[0057] Figure 11cMicrostructure diagram of 0RE experimental steel in the embodiment of the application after 15 min ART annealing at 660 DEG C;

[0058] Figure 11d Microstructure diagram of 48ppmRE experimental steel in the embodiment of the application after 15 min ART annealing at 620 DEG C;

[0059] Figure 11e Microstructure diagram of 48ppmRE experimental steel in the embodiment of the application after 15 min ART annealing at 640 DEG C;

[0060] Figure 11f Microstructure diagram of 48ppmRE experimental steel in the embodiment of the application after 15 min ART annealing at 660 DEG C;

[0061] Figure 12a TEM image of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0062] Figure 12b TEM image of 48ppmRE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0063] Figure 13a Bulk austenite TEM image of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0064] Figure 13b Diffraction spot of bulk austenite of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0065] Figure 13c Lath austenite TEM image of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0066] Figure 13d Diffraction spot of lath austenite of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0067] Figure 14a M / A island structure of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0068] Figure 14b Austenite diffraction spot of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0069] Figure 14c Martensite diffraction spot of 0RE experimental steel in the embodiment of the application after 640 DEG C ART annealing;

[0070] Figure 15a XRD phase analysis diagram of 0RE experimental steel in the embodiment of the application under different ART annealing temperatures;

[0071] Figure 15b XRD phase analysis diagram of the 48ppm RE experimental steel at different ART annealing temperatures in the embodiment of the application;

[0072] Figure 15c Graph of austenite volume fraction of the experimental steel at different ART annealing temperatures in the embodiment of the application;

[0073] Figure 16a Stress-strain curve diagram of the 0 RE experimental steel at different ART annealing temperatures in the embodiment of the application;

[0074] Figure 16b Stress-strain curve diagram of the 48ppm RE experimental steel at different ART annealing temperatures in the embodiment of the application;

[0075] Figure 16c Yield strength curve diagram of the experimental steel at different ART annealing temperatures in the embodiment of the application;

[0076] Figure 16d Tensile strength curve diagram of the experimental steel at different ART annealing temperatures in the embodiment of the application;

[0077] Figure 16e Elongation curve diagram of the experimental steel at different ART annealing temperatures in the embodiment of the application;

[0078] Figure 16f Product of strength and ductility curve diagram of the experimental steel at different ART annealing temperatures in the embodiment of the application;

[0079] Figure 17a Fracture morphology diagram of the 0 RE experimental steel at 620 DEG C ART annealing in the embodiment of the application;

[0080] Figure 17b Fracture morphology diagram of the 0 RE experimental steel at 640 DEG C ART annealing in the embodiment of the application;

[0081] Figure 17c Fracture morphology diagram of the 0 RE experimental steel at 660 DEG C ART annealing in the embodiment of the application;

[0082] Figure 17d Fracture morphology diagram of the 48ppm RE experimental steel at 620 DEG C ART annealing in the embodiment of the application;

[0083] Figure 17e Fracture morphology diagram of the 48ppm RE experimental steel at 640 DEG C ART annealing in the embodiment of the application;

[0084] Figure 17f Fracture morphology diagram of the 48ppm RE experimental steel at 660 DEG C ART annealing in the embodiment of the application;

[0085] Figure 18aMicrostructure diagram of 0RE experimental steel ART annealed at 640 DEG C for 5 min in the embodiment of the present application;

[0086] Figure 18b Microstructure diagram of 0RE experimental steel ART annealed at 640 DEG C for 15 min in the embodiment of the present application;

[0087] Figure 18c Microstructure diagram of 0RE experimental steel ART annealed at 640 DEG C for 25 min in the embodiment of the present application;

[0088] Figure 18d Microstructure diagram of 48ppmRE experimental steel ART annealed at 640 DEG C for 5 min in the embodiment of the present application;

[0089] Figure 18e Microstructure diagram of 48ppmRE experimental steel ART annealed at 640 DEG C for 15 min in the embodiment of the present application;

[0090] Figure 18f Microstructure diagram of 48ppmRE experimental steel ART annealed at 640 DEG C for 25 min in the embodiment of the present application;

[0091] Figure 19a XRD phase analysis diagram of 0RE experimental steel under different ART annealing time in the embodiment of the present application;

[0092] Figure 19b XRD phase analysis diagram of 48ppmRE experimental steel under different ART annealing time in the embodiment of the present application;

[0093] Figure 19c Gurley volume fraction diagram of experimental steel under different ART annealing time in the embodiment of the present application;

[0094] Figure 20a Stress-strain curve diagram of 0RE experimental steel under different ART annealing time in the embodiment of the present application;

[0095] Figure 20b Stress-strain curve diagram of 48ppmRE experimental steel under different ART annealing time in the embodiment of the present application;

[0096] Figure 20c Yield strength curve diagram of experimental steel under different ART annealing time in the embodiment of the present application;

[0097] Figure 20d Tensile strength curve diagram of experimental steel under different ART annealing time in the embodiment of the present application;

[0098] Figure 20e Elongation curve diagram of experimental steel under different ART annealing time in the embodiment of the present application;

[0099] Figure 20fThe steel strength-plasticity product curve under different ART annealing time in the embodiment of the application is drawn;

[0100] Figure 21a The fracture morphology graph of the 0RE experimental steel ART annealed at 640 DEG C for 5 min in the embodiment of the application is drawn;

[0101] Figure 21b The fracture morphology graph of the 0RE experimental steel ART annealed at 640 DEG C for 15 min in the embodiment of the application is drawn;

[0102] Figure 21c The fracture morphology graph of the 0RE experimental steel ART annealed at 640 DEG C for 25 min in the embodiment of the application is drawn;

[0103] Figure 21d The fracture morphology graph of the 48ppmRE experimental steel ART annealed at 640 DEG C for 5 min in the embodiment of the application is drawn;

[0104] Figure 21e The fracture morphology graph of the 48ppmRE experimental steel ART annealed at 640 DEG C for 15 min in the embodiment of the application is drawn;

[0105] Figure 21f The fracture morphology graph of the 48ppmRE experimental steel ART annealed at 640 DEG C for 25 min in the embodiment of the application is drawn;

[0106] Figure 22a The microstructure graph of the 0RE experimental steel directly annealed at 620 DEG C for 15 min in the embodiment of the application is drawn;

[0107] Figure 22b The microstructure graph of the 0RE experimental steel directly annealed at 640 DEG C for 15 min in the embodiment of the application is drawn;

[0108] Figure 22c The microstructure graph of the 0RE experimental steel directly annealed at 660 DEG C for 15 min in the embodiment of the application is drawn;

[0109] Figure 22d The microstructure graph of the 48ppmRE experimental steel directly annealed at 620 DEG C for 15 min in the embodiment of the application is drawn;

[0110] Figure 22e The microstructure graph of the 48ppmRE experimental steel directly annealed at 640 DEG C for 15 min in the embodiment of the application is drawn;

[0111] Figure 22f The microstructure graph of the 48ppmRE experimental steel directly annealed at 660 DEG C for 15 min in the embodiment of the application is drawn;

[0112] Figure 23a The TEM image of the 0RE experimental steel directly annealed at 640 DEG C in the embodiment of the application is drawn;

[0113] Figure 23b TEM image of 48ppm RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0114] Figure 23c Dislocation of 0 RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0115] Figure 24a TEM image of austenite of 0 RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0116] Figure 24b Diffraction spot of austenite of 0 RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0117] Figure 24c TEM image of austenite of 48ppm RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0118] Figure 24d Diffraction spot of austenite of 48ppm RE experimental steel after direct annealing at 640 DEG C in two-phase region in the embodiment of the present application;

[0119] Figure 25a XRD phase analysis diagram of 0 RE experimental steel at different direct annealing temperatures in two-phase region in the embodiment of the present application;

[0120] Figure 25b XRD phase analysis diagram of 48ppm RE experimental steel at different direct annealing temperatures in two-phase region in the embodiment of the present application;

[0121] Figure 25c Austenite volume fraction diagram of experimental steel at different direct annealing temperatures in two-phase region in the embodiment of the present application;

[0122] Figure 26a Stress-strain curve diagram of 0 RE experimental steel at different critical annealing temperatures after direct annealing in two-phase region in the embodiment of the present application;

[0123] Figure 26b Stress-strain curve diagram of 48ppm RE experimental steel at different critical annealing temperatures after direct annealing in two-phase region in the embodiment of the present application;

[0124] Figure 26c Yield strength curve diagram of experimental steel at different critical annealing temperatures after direct annealing in two-phase region in the embodiment of the present application;

[0125] Figure 26d Tensile strength curve diagram of experimental steel at different critical annealing temperatures after direct annealing in two-phase region in the embodiment of the present application;

[0126] Figure 26eThe elongation rate curve diagram of the experimental steel under different critical annealing temperatures in two-phase zone direct annealing in the embodiment of the application is shown in the figure;

[0127] Figure 26f The strength-plasticity product curve diagram of the experimental steel under different critical annealing temperatures in two-phase zone direct annealing in the embodiment of the application is shown in the figure;

[0128] Figure 27a The fracture morphology diagram of the 0RE experimental steel under two-phase zone direct annealing at 620 DEG C in the embodiment of the application is shown in the figure;

[0129] Figure 27b The fracture morphology diagram of the 0RE experimental steel under two-phase zone direct annealing at 640 DEG C in the embodiment of the application is shown in the figure;

[0130] Figure 27c The fracture morphology diagram of the 0RE experimental steel under two-phase zone direct annealing at 660 DEG C in the embodiment of the application is shown in the figure;

[0131] Figure 27d The fracture morphology diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 620 DEG C in the embodiment of the application is shown in the figure;

[0132] Figure 27e The fracture morphology diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 640 DEG C in the embodiment of the application is shown in the figure;

[0133] Figure 27f The fracture morphology diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 660 DEG C in the embodiment of the application is shown in the figure;

[0134] Figure 28a The microstructure diagram of the 0RE experimental steel under two-phase zone direct annealing at 640 DEG C for 5 min in the embodiment of the application is shown in the figure;

[0135] Figure 28b The microstructure diagram of the 0RE experimental steel under two-phase zone direct annealing at 640 DEG C for 15 min in the embodiment of the application is shown in the figure;

[0136] Figure 28c The microstructure diagram of the 0RE experimental steel under two-phase zone direct annealing at 640 DEG C for 25 min in the embodiment of the application is shown in the figure;

[0137] Figure 28d The microstructure diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 640 DEG C for 5 min in the embodiment of the application is shown in the figure;

[0138] Figure 28e The microstructure diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 640 DEG C for 15 min in the embodiment of the application is shown in the figure;

[0139] Figure 28f The microstructure diagram of the 48ppmRE experimental steel under two-phase zone direct annealing at 640 DEG C for 25 min in the embodiment of the application is shown in the figure;

[0140] Figure 29a XRD phase analysis diagram of 0RE experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0141] Figure 29b XRD phase analysis diagram of 48ppmRE experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0142] Figure 29c Graph of austenite volume fraction of experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0143] Figure 30a Stress-strain curve diagram of 0RE experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0144] Figure 30b Stress-strain curve diagram of 48ppmRE experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0145] Figure 30c Yield strength curve diagram of experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0146] Figure 30d Tensile strength curve diagram of experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0147] Figure 30e Elongation curve diagram of experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0148] Figure 30f Product of strength and ductility curve diagram of experimental steel in different two-phase regions under direct annealing time in the embodiment of the application;

[0149] Figure 31a Fracture morphology diagram of 0RE experimental steel after direct annealing in a two-phase region at 640 DEG C for 5 min in the embodiment of the application;

[0150] Figure 31b Fracture morphology diagram of 0RE experimental steel after direct annealing in a two-phase region at 640 DEG C for 15 min in the embodiment of the application;

[0151] Figure 31c Fracture morphology diagram of 0RE experimental steel after direct annealing in a two-phase region at 640 DEG C for 25 min in the embodiment of the application;

[0152] Figure 31d Fracture morphology diagram of 48ppmRE experimental steel after direct annealing in a two-phase region at 640 DEG C for 5 min in the embodiment of the application;

[0153] Figure 31eFracture morphology diagram of 48ppmRE experimental steel after direct annealing in two-phase region at 640 DEG C for 15 min in the embodiment of the application;

[0154] Figure 31f Fracture morphology diagram of 48ppmRE experimental steel after direct annealing in two-phase region at 640 DEG C for 25 min in the embodiment of the application;

[0155] Figure 32a SEM diagram of 0RE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0156] Figure 32b SEM diagram of 48ppmRE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0157] Figure 32c SEM diagram of 0RE experimental steel under the best condition of ART annealing in the application;

[0158] Figure 32d SEM diagram of 48ppmRE experimental steel under the best condition of ART annealing in the application;

[0159] Figure 33a TEM diagram a of 0RE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0160] Figure 33b TEM diagram a of 48ppmRE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0161] Figure 33c TEM diagram a of 0RE experimental steel under the best condition of ART annealing in the application;

[0162] Figure 33d TEM diagram a of 48ppmRE experimental steel under the best condition of ART annealing in the application;

[0163] Figure 34a TEM diagram b of 0RE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0164] Figure 34b TEM diagram b of 48ppmRE experimental steel under the best condition of direct annealing in two-phase region in the application;

[0165] Figure 34c TEM diagram b of 0RE experimental steel under the best condition of ART annealing in the application;

[0166] Figure 34d TEM diagram b of 48ppmRE experimental steel under the best condition of ART annealing in the application;

[0167] Figure 35A plot of the plastic deformation resistance of the crystal versus the dislocation density;

[0168] Figure 36a A plot of the mechanical properties of the 0RE experimental steel under two annealing processes in the application;

[0169] Figure 36b A plot of the mechanical properties of the 48ppmRE experimental steel under two annealing processes in the application;

[0170] Figure 37a An ODF section view of the 0RE experimental steel under ART annealing in the application, φ2=45°;

[0171] Figure 37b An ODF section view of the 48ppmRE experimental steel under ART annealing in the application, φ2=45°;

[0172] Figure 37c An ODF section view of the 0RE experimental steel under two-phase zone direct annealing in the application, φ2=45°;

[0173] Figure 37d An ODF section view of the 48ppmRE experimental steel under two-phase zone direct annealing in the application, φ2=45°;

[0174] Figure 38a A comparison plot of the texture content of the two experimental steels under ART annealing in the application;

[0175] Figure 38b A comparison plot of the texture content of the two experimental steels under two-phase zone direct annealing in the application;

[0176] Figure 39a A microstructure view of the 48ppmRE experimental steel with a deformation degree of 73% in the application;

[0177] Figure 39b A microstructure view of the 48ppmRE experimental steel with a deformation degree of 50% in the application;

[0178] Figure 39c A microstructure view of the 48ppmRE experimental steel with a deformation degree of 33% in the application;

[0179] Figure 40a A texture view of the 48ppmRE experimental steel with a deformation degree of 33% in the application;

[0180] Figure 40b A texture view of the 48ppmRE experimental steel with a deformation degree of 50% in the application;

[0181] Figure 40c A texture view of the 48ppmRE experimental steel with a deformation degree of 73% in the application;

[0182] Figure 41a ODF section of 48ppmRE experimental steel with 33% deformation in the present application φ2 = 45°;

[0183] Figure 41b ODF section of 48ppmRE experimental steel with 50% deformation in the present application;

[0184] Figure 41c ODF section of 48ppmRE experimental steel with 73% deformation in the present application;

[0185] Figure 42 Texture content column chart of 48ppmRE experimental steel with different deformations in the present application;

[0186] Figure 43 Pressure-strain curve chart of 48ppmRE experimental steel with different deformations in the present application;

[0187] Figure 44 Technical route schematic diagram of the present application;

[0188] Figure 45 Cold bending state schematic diagram of 48ppmRE experimental steel with 73% deformation in the present application after two-phase region direct annealing. DETAILED DESCRIPTION

[0189] 1.1 Preparation of experimental materials

[0190] 1.1.1 Composition design

[0191] Super-high-strength automobile steel has excellent mechanical properties and forming properties, which benefits from obtaining a proper amount of metastable residual austenite at room temperature. The volume fraction of residual austenite can be controlled by the distribution of C and Mn, and other alloying elements also play a key role in the stability of austenite. In this embodiment, 0.1C-5Mn medium manganese steel is taken as the basis, and the chemical compositions of two groups of experimental steels are designed, one of which adds 48ppm Ce, and the other does not add rare earth as a contrast sample. The content of each element of the experimental steel is shown in Table 7. The role of rare earth element Ce is to improve the plasticity and toughness of the steel, and its mechanism is to improve the toughness of the steel by changing the morphology and distribution of inclusions, and to improve the delayed cracking. In addition, the addition of rare earth elements can reduce the stacking fault energy of austenite, increase the martensite transformation point, and thus increase the transformation amount of austenite, and improve the plasticity and toughness of the steel.

[0192] Table 7 Chemical composition of experimental steel (mass fraction, %)

[0193]

[0194] 1.1.2 Research route

[0195] In this embodiment, the experimental steel is sequentially subjected to smelting and casting, hot rolling, pickling, cold rolling and heat treatment. The various properties of the experimental steel after heat treatment are tested, and the data obtained by testing are analyzed and sorted, as shown in Table 1. Figure 44 As shown in Table 1, three deformation degrees are set in the cold rolling process, which are 33%, 50% and 73%. The reason for selecting these three deformation degrees is to simulate the typical deformation degrees of the commonly used thickness specifications of the super-high-strength automobile steel after cold rolling when the automobile parts are processed. The heat treatment processes include two-phase region direct annealing and ART annealing, which are the most mainstream heat treatment processes for super-high-strength automobile steel.

[0196] (1) Smelting and casting

[0197] A 50 kg multifunctional vacuum induction furnace is selected for smelting. Before smelting, 20 min of vacuum is extracted, and the vacuum degree in the furnace is controlled below 70 Pa. Then, the temperature is heated to 1600℃, and the heating time is 90 min. After the raw materials are completely melted, the alloy elements are added. After the alloy elements are melted, 5 min of homogenization treatment is performed, and then casting is performed.

[0198] (2) Hot rolling

[0199] The Ф450 double-roller hot rolling machine is used for hot rolling. The casted blank is heated to 1200℃ for 5 h for diffusion annealing and then rolled. The opening rolling temperature is 1170-1180℃, and the final rolling temperature is 860℃. The specific rolling process is shown in Table 8.

[0200] Table 8 Distribution of rolling amount of experimental steel

[0201]

[0202] (3) Pickling

[0203] After hot rolling, the plate has oxidation skin. If direct cold rolling is performed, the surface quality of the steel plate will be poor, and the subsequent coating performance and spraying performance will also be affected. Therefore, the oxidation skin on the surface of the hot-rolled plate needs to be removed before cold rolling. The method for removing the oxidation skin is usually chemical method, i.e. pickling. The hot-rolled plate is soaked in concentrated hydrochloric acid with pH of 0 for 10 h. After the oxidation skin is removed, the hydrochloric acid is washed off with clean water, and then the plate is blown dry with a hair dryer and coated with oil to prevent oxidation.

[0204] (4) Cold rolling

[0205] The cold rolling is carried out by using a vertical four-high reversible cold rolling mill, and the maximum rolling force is 2500 KN. Since the strength of the steel plate is too high, annealing softening is carried out before rolling, and the annealing temperature is 850 DEG C and the time is 4 hours. The annealed hot-rolled plate is divided into three groups, and is rolled to 2 mm, 1.5 mm and 0.8 mm respectively, and the deformation degrees are 33%, 50% and 73% respectively. The purpose of annealing softening is to prevent the steel plate from being too high in strength and being unable to be successfully rolled to the target thickness by the rolling mill; and the annealing softening temperature is too high or the annealing softening time is too long, which easily leads to the re-formation of oxide scale on the steel plate and the grain coarsening.

[0206] (5) Two-phase region direct annealing / ART annealing

[0207] The CCT curve of the experimental steel is obtained by simulation by using JMatPro software. Figure 2a The simulation results of the 0RE experimental steel are shown in Table 1, Figure 2b The simulation results of the 48ppm experimental steel are shown in Table 2. The Ac1 and Ac3 of the experimental steel without rare earth (0RE) are 609 DEG C and 730 DEG C respectively, and the Ac1 and Ac3 of the experimental steel with rare earth (48ppm) are 598 DEG C and 736 DEG C respectively. The DSC differential thermal analyzer is used for testing, the highest temperature is set to 900 DEG C, the heating rate is 10 DEG C / min, and it is measured that the Ac1 and Ac3 of the experimental steel without rare earth are 594 DEG C and 736 DEG C respectively, and the Ac1 and Ac3 of the experimental steel with rare earth are 598 DEG C and 748 DEG C respectively.

[0208] The ultra-high strength automobile steel can obtain a large number of metastable reverse phase transformation austenite by the method of heat treatment, and the retention of austenite makes the TRIP effect continue to occur when deformation occurs, thereby obtaining a high-strength and high-plasticity structure, and improving the mechanical properties and stamping performance. In the embodiment, two-phase region direct annealing and ART annealing are used as the heat treatment processes, and the heat treatment temperature is determined according to the simulation and the two-phase region temperature range determined by DSC measurement.

[0209] The specific heat treatment process is that the ART annealing is to first quench the cold-rolled plate after holding at 800 DEG C for 5 min (water quenching). The quenching temperature is obtained by multiple experiments, and the quenching temperature is too high and the holding time is too long, which will make the austenite grain grow, and affect the performance of the material. After quenching, annealing treatment is carried out, and the annealing process parameters are as follows: 620 DEG C-15 min, 640 DEG C-15 min, 660 DEG C-15 min, and 640 DEG C-5 min, 640 DEG C-15 min, 640 DEG C-25 min, and air cooling to room temperature; the two-phase region direct annealing process does not quench, but directly carries out annealing treatment, and the annealing process parameters are the same as those of the ART process. The ART annealing process is as shown in Figure 3a The two-phase region direct annealing process is as shown in Figure 3b ​The grain size distribution of the experimental steel with rare earth Ce after ART annealing is more uniform, and the comprehensive mechanical properties are better. Compared with the ART process, the two-phase region direct annealing process saves the quenching process before annealing, greatly shortens the time, and the microstructure and properties of the experimental steel after two-phase region direct annealing are similar to or even better than those of the ART process.

[0210] 1.2 Preparation and characterization method of experimental material

[0211] 1.2.1 Tensile test at room temperature

[0212] The cold-rolled plate was cut into proportional samples along the rolling direction using a wire cutting machine, and the sample size was as shown in Figure 4 The original gauge length was calculated by formula (1-1), and the tensile strength, yield strength and elongation after fracture of the sample were tested using SHT-4605 type microcomputer controlled electronic universal testing machine.

[0213]

[0214] 1.2.2 Scanning electron microscope (SEM) test

[0215] Phenom LE field emission scanning electron microscope provided by Phenom Company of the Netherlands was used to observe the microstructure morphology.

[0216] 1.2.3 Transmission electron microscope (TEM) test

[0217] JEM-2010 type transmission electron microscope was used to analyze the microstructure morphology and dislocation in the experimental steel, and the composition of the phase was determined by electron diffraction.

[0218] 1.2.4 Observation of original austenite

[0219] The experimental steel was quenched at 800℃ for 5min, and after grinding, polishing and etching of the quenched sample, the original austenite microstructure was observed by metallographic microscope, and the original austenite grain size distribution of the two experimental steels was measured using Nano Measurer software.

[0220] 1.2.5 X-ray diffraction method (XRD) test

[0221] X Pert PRO MPD type X-ray diffractometer of Panalytical Company of the Netherlands was used to measure the austenite content of the experimental steel, and formula (1-2) was used to calculate the volume fraction of residual austenite:

[0222]

[0223] Wherein G is the crystal face parameter related value corresponding to austenite phase and martensite phase, different G values are shown in Table 9. The volume fraction of austenite VA corresponding to different IM / IA is calculated respectively, and then the arithmetic mean value is calculated, which is the volume fraction of residual austenite.

[0224] The X-ray diffractometer is used to test the texture of the sample, the ODF cross section is drawn by Roe method, and the volume fraction of each texture component is calculated by ResMat-TexTools texture analysis software.

[0225] Table 9 G values corresponding to different crystal face indices

[0226]

[0227] 2. Effect of Ce on microstructure and properties of ultra-high strength automobile steel under ART process

[0228] The addition of rare earth not only can purify the molten steel and change the morphology of inclusions, but also can play a role in solid solution strengthening, so as to improve the mechanical properties and stamping properties of the steel. In this part, two groups of experimental steels with and without rare earth will be selected, and the two groups of experimental steels will be annealed by ART, and then the effect of rare earth Ce on the microstructure and properties of experimental steel under ART process will be explored through mechanical property test and microstructure morphology analysis.

[0229] 2.1 Analysis of original microstructure and properties

[0230] 2.1.1 Original austenite grain

[0231] Figure 5a 、 Figure 5b Effect of rare earth elements on original austenite morphology and grain size. As can be seen from the figure, the austenite grain of the experimental steel with rare earth element is smaller. This is because the addition of rare earth element increases the phase transition point of the experimental steel, and the experimental steel without rare earth element nucleates before the experimental steel with rare earth element when heated to austenitizing temperature. Another reason is that rare earth elements not only can be adsorbed on the grain boundary, but also can form inclusions, which will affect the growth direction and speed of the grain, making the grain growth more complex, which helps to refine the grain. The average grain size of the austenite of the two groups of experimental steels is measured by NanoMeasurer software, which is 4.91 μm and 4.42 μm respectively, and the grain size distribution graph is shown in Figure 5c 、 Figure 5d The grain size of the experimental steel without rare earth is less than 3 μm, accounting for 9.75%, and the grain size between 3-5 μm accounts for 48%. While the grain size of the experimental steel with rare earth is less than 3 μm, accounting for 18.75%, and the grain size between 3-5 μm accounts for 49%. Through comparison, it can be known that the grain size of the experimental steel with rare earth is smaller.

[0232] 2.1.2 Microstructure and mechanical properties of hot-rolled plate

[0233] Figure 6a and Figure 6b is the microstructure of the hot-rolled plate, Figure 6a is the microstructure of the hot-rolled plate without rare earth, Figure 6b is the microstructure of the hot-rolled plate containing 48ppm Ce. As can be seen from the figure, the hot-rolled plate is lath martensite and ferrite, which is due to the elongation of the grains during rolling and the retention of the transformation of austenite to martensite during laminar cooling.

[0234] Figure 7 is the stress-strain curve of the two hot-rolled plates. As can be seen from the figure, the tensile strength of the experimental steel without rare earth is 1570MPa, the yield strength is 1310MPa, the elongation is 8.74%, and the product of strength and plasticity is 13.7GPa·%, while the tensile strength of the experimental steel containing rare earth is 1620MPa, the yield strength is 1250MPa, the elongation is 9.95%, and the product of strength and plasticity is 16.2GPa·%. The comprehensive mechanical properties are better than those of the experimental steel without rare earth. Due to the existence of a large amount of martensite and deformed ferrite in the microstructure of the hot-rolled plate, and a large amount of dislocations in the ferrite, the strength of the hot-rolled plate is high and the plasticity is poor due to dislocation strengthening and martensite strengthening.

[0235] 2.1.3 Microstructure and properties of cold-rolled plate

[0236] Figure 8a , Figure 8b is the original microstructure of the cold-rolled plate without heat treatment. Since the plate has been rolled, there are both elongated lath ferrite and martensite in the microstructure, as well as blocky ferrite and martensite, which is significantly different from the single martensite structure after quenching. By comparing the microstructure of the cold-rolled plates of the two experimental steels, it is found that the grains of the experimental steel containing rare earth are finer, indicating that the addition of rare earth can refine the grains and improve the material properties.

[0237] Figure 9 is the stress-strain curve of the cold-rolled plate. As can be seen from the figure, the tensile strength of the experimental steel without rare earth and the experimental steel containing rare earth is 1764MPa and 1850MPa respectively, the elongation is 6.9% and 7.5% respectively, and the product of strength and plasticity is 12.2GPa·% and 13.9GPa·% respectively, indicating that the addition of rare earth improves the mechanical properties of the material. Due to the high distortion energy of the martensite and deformed ferrite in the microstructure of the cold-rolled plate, the strength of the cold-rolled plate is very high and the elongation is very low, so further heat treatment is needed to improve the mechanical properties and stamping properties.

[0238] 2.2 Effect of ART process on microstructure and properties of rare earth microalloyed ultra-high strength automobile steel

[0239] 2.2.1 Quenching microstructure of cold-rolled plate

[0240] SEM microstructure of the cold-rolled plate after quenching is shown in Figure 10a and Figure 10b . Figure 10a The 0RE experimental steel without rare earth, Figure 10b The experimental steel containing 48ppm Ce, as can be seen from the figure, the microstructure of the experimental steel after quenching is lath martensite, and C precipitates around the lath martensite. Moreover, the martensite after quenching is formed inside the original austenite grain, still retaining the grain boundary of the original austenite. Since the martensite is a hard phase structure, the strength of the quenched material is very high and the plasticity is very poor, which needs further annealing treatment.

[0241] 2.2.2 Effect of annealing temperature on microstructure and properties

[0242] ART annealing treatment was performed on the two groups of experimental steels. The steel was first heated to a temperature above Ac3 for quenching treatment. Quenching temperature that is too high and holding time that is too long will cause the growth of austenite grains, affecting the performance of the material. After repeated experiments, the quenching temperature was determined to be 800℃ and the holding time was 5min. Figure 11a- Figure 11f are microstructure photos of the two groups of experimental steels after holding for 15min at different annealing temperatures. As can be seen from the figure, the microstructure after annealing is a mixture of austenite, newly formed martensite and ferrite. During quenching, martensite nucleates inside the original austenite grain, and austenite also forms at the original austenite grain boundary. During annealing, due to the redistribution of C and Mn elements, the C element in part of the martensite becomes unsaturated and transforms into ferrite.

[0243] As shown in Figure 11a- Figure 11f , the protruding part is newly formed martensite and austenite, and the concave part is ferrite. When the annealing temperature is 620℃, as shown in Figure 11a , Figure 11d , the protruding structure is austenite and martensite, and the concave part is ferrite, and the carbide in the quenched state is basically not dissolved. When the annealing temperature is increased to 640℃, as shown in Figure 11b , Figure 11e , at this time the carbide begins to dissolve, providing energy for the nucleation of austenite, and the content of austenite increases significantly. When the annealing temperature is increased to 660℃, as shown in Figure 11c , Figure 11fThe microstructure is mainly austenite and ferrite, which is due to the temperature rise, carbide is basically dissolved, providing sufficient energy for the nucleation of austenite, but due to the increase of temperature, the grain size is also large, so that the C, Mn elements in austenite are reduced, and the austenite becomes unstable, and part of it is transformed into martensite during cooling. By comparing the two groups of experimental steels, it is found that the grain of the experimental steel containing rare earth is smaller, because the atomic radius of rare earth is larger than that of Fe, and can only be dissolved in the defects in the crystal, and cannot be dissolved in austenite, and the defects at the grain boundary are more, so the rare earth is gathered at the grain boundary, hindering the atomic diffusion, thereby inhibiting the grain growth. The rare earth atoms also have the effect of desulfurization and deoxidization, forming rare earth compounds in the matrix, which can also act as heterogeneous nuclei for grain refinement.

[0244] Figure 12a 、 Figure 12b The TEM images of the two groups of experimental steels after annealing at 640℃ are shown in the figure, it can be seen from the figure that the microstructure after annealing is composed of ferrite, lath-shaped austenite and martensite, and dislocations exist in the ferrite. These dislocations are formed by the phase transition from face-centered cubic austenite to body-centered cubic ferrite at high temperature, while dislocations rarely appear in the interior of martensite, because the atomic arrangement of martensite is very tight, and it is difficult to appear dislocations, which is also the reason for the high strength of martensite.

[0245] Figure 13a- Figure 13d TEM images of different austenite morphologies formed after the experimental steel is annealed at 640℃, Figure 13a 、 Figure 13b TEM images and diffraction spots of blocky austenite, Figure 13c 、 Figure 13d TEM images and diffraction spots of lath-shaped austenite, blocky austenite is formed by the meeting of lath-shaped austenite with the same orientation. Since austenite is a high-temperature phase, it is necessary to add appropriate alloying elements and change the heat treatment process parameters to retain it at room temperature. When the annealing time is 640℃, the stability of austenite is the highest, and the content of residual austenite reaches 25.7%. Figure 14a- Figure 14c M / A island structure and diffraction spots retained after the 0RE experimental steel is annealed at 640℃, Figure 14b 、 Figure 14c Diffraction spots of austenite and martensite, some unstable austenite will be transformed into martensite during cooling, and a part of austenite does not completely transform into M / A island structure.

[0246] The two groups of experimental steels after ART annealing are tested by XRD, and the phase analysis diagram and the volume fraction of austenite are shown in Figure 15a- Figure 15c , Figure 15a 、 Figure 15bThe phase analysis diagrams of the experimental steels without and with rare earth, respectively, wherein γ represents the FCC phase, a represents the BCC phase, and the austenite volume fraction can be calculated from the diffraction peak intensity of austenite and ferrite. It can be seen from Figure 15a 、 Figure 15b that the austenite diffraction peak gradually increases with the increase of the annealing temperature and reaches the maximum at 640°C, indicating that the austenite content is the highest at 800°C / 640°C-15 min. When the temperature rises to 660°C, the FCC diffraction peak begins to decrease, and the BCC diffraction peak increases, which shows that after the annealing temperature exceeds 640°C, the austenite volume fraction decreases with the increase of the annealing temperature, and new martensite is generated. By comparing the XRD diagrams of the two groups of experimental steels, it is found that the FCC diffraction peak of the experimental steel containing rare earth is stronger than that of the experimental steel without rare earth, indicating that the rare earth element is very helpful to the increase of the austenite volume fraction. Figure 15c The austenite volume fraction calculated is the austenite volume fraction of the two groups of experimental steels, which is first increased and then decreased with the increase of the annealing temperature, and reaches the maximum at 640°C. The austenite volume fraction of the experimental steels without and with rare earth reaches 24.7% and 25.7%, respectively. The reason why the austenite volume fraction decreases with the increase of the annealing temperature is that the austenite grains are coarsened, and part of the austenite is transformed into martensite during the cooling process.

[0247] Figure 16a- Figure 16f The stress-strain curves and mechanical property diagrams of the two groups of experimental steels under different annealing temperatures by ART can be found that the yield strength, tensile strength, elongation and strength-plasticity product of the experimental steel with the addition of rare earth element Ce are improved compared with the experimental steel without rare earth.

[0248] As shown in Figure 16c , the yield strength gradually decreases with the increase of the temperature, and reaches the maximum at 620°C, which is 959 MPa and 954 MPa for the experimental steels without and with rare earth, respectively. The yield strength is mainly affected by the grain size and dislocation. With the increase of the annealing temperature, the grain size increases, so the yield strength gradually decreases. As shown in Figure 16d , the tensile strength gradually increases with the increase of the temperature, which is due to the fact that with the increase of the annealing temperature, part of the austenite is transformed into new martensite during the cooling process, and the TRIP effect is also affected during the tensile process of the sample, so the tensile strength gradually increases with the increase of the temperature. The tensile strength has a maximum value at 660°C, which is 1289 MPa and 1175 MPa, respectively.

[0249] As shown in Figure 16eAs shown in the figure, the elongation first increases and then decreases with the increase of temperature, and reaches the maximum value at 640℃ annealing. The elongation of the experimental steel without rare earth is 33.1%, and the elongation of the experimental steel with rare earth element Ce reaches 34.8%. The elongation is affected by TRIP effect. When TRIP effect occurs, unstable austenite is transformed into martensite, which delays the occurrence of necking phenomenon. Therefore, the higher the austenite content, the more significant the TRIP effect, and the higher the elongation. This is consistent with the analysis results of the austenite content. The product of strength and plasticity first increases and then decreases with the increase of temperature, and the mechanical properties reach the best at 800℃-5min / 640℃-15min. The product of strength and plasticity of the two experimental steels is 33.1GPa·% and 34.8GPa·% respectively.

[0250] Figure 17a- Figure 17f The tensile fracture morphology of the two experimental steels after ART at different temperatures. The fracture of the material can be divided into brittle fracture and plastic fracture, and the fracture morphology after the two kinds of fracture is very different. The fracture of brittle fracture is usually cleavage fracture, and the morphology of cleavage fracture is river-like, and almost no dimple is produced. Plastic fracture often appears fibrous dimple, and the proportion of dimple can be used to evaluate the plasticity of the material.

[0251] As shown in the figure, Figure 17a , Figure 17d When the annealing temperature is 620℃, the fracture is mainly cleavage fracture, and there are few dimples. When the temperature rises to 640℃, as shown in the figure, Figure 17b , Figure 17e The dimples are densely distributed and small in size, which belongs to ductile fracture. When the annealing temperature reaches 660℃, as shown in the figure, Figure 17c , Figure 17f The number of dimples is significantly reduced, and the size is larger, which also belongs to ductile fracture. The post-fracture elongation of steel is related to the number, size and density of dimples. By comparing the fracture morphology at three different temperatures, it is found that when the annealing temperature is 640℃, the dimples are the smallest, the number is the most and the distribution is uniform, and the plasticity is the best, which is consistent with the results reflected by the stress-strain curve. Figure 17a , Figure 17b , Figure 17c The fracture morphology of the experimental steel without rare earth is compared with that of the experimental steel containing rare earth, Figure 17d , Figure 17e , Figure 17f It is found that the dimples of the experimental steel containing rare earth are more uniformly distributed and more in number, which shows that the addition of rare earth can improve the plasticity and toughness of the material. Studies have shown that ductile fracture is mainly caused by austenite deformation, and cleavage fracture is related to the deformation of ferrite, which is consistent with the XRD results. When the annealing temperature is 640℃, the content of residual austenite is the highest, and the content of austenite in the rare earth steel under the same heat treatment process is higher than that in the steel without rare earth.

[0252] 2.2.3 Effect of annealing time on microstructure and properties

[0253] Two groups of experimental steels were quenched at 800℃ for 5 minutes, and then annealed at 640℃ for 5 minutes, 15 minutes, and 25 minutes respectively. The microstructures after air cooling to room temperature were as follows: Figure 18a- Figure 18f As shown.

[0254] like Figure 18a , Figure 18d As shown, when the annealing time is 5 minutes, the microstructure mainly consists of ferrite and austenite. The dark-colored, concave microstructure is ferrite, and the light-colored, convex microstructure is austenite. A large amount of carbides are also dispersed within the ferrite. This is because the holding time is short, and the carbides do not have enough time to dissolve. At this point, the austenite nucleation does not have sufficient energy, resulting in a low content of retained austenite. Figure 18b , Figure 18e As shown, the holding time at this point is 15 minutes. The microstructure is still mainly composed of ferrite and austenite. Due to the extended holding time, the carbides have sufficient time to dissolve, providing enough energy for austenite nucleation. The volume fraction of austenite retained at room temperature is the highest. Figure 18c , Figure 18f As shown, the annealing time was 25 minutes. At this time, the microstructure contained not only ferrite and austenite, but also newly formed martensite. This is because the extended holding time allowed the carbides to largely dissolve, giving the austenite sufficient time to nucleate and grow. The grown austenite became unstable and partially transformed into martensite during cooling. By comparing the two groups of experimental steels, it was found that the microstructure of the rare earth-containing experimental steel was finer and more uniform, which may be related to the microalloying of rare earth elements, resulting in better mechanical and stamping properties.

[0255] X-ray diffraction (XRD) was used to analyze the austenite content of two groups of experimental steel samples annealed for different ART times to clarify the effect of annealing time on austenite content and microstructure. The results are as follows: Figure 19a Figure 19c As shown.

[0256] In the figure, α represents body-centered cubic (BCC), γ represents face-centered cubic (FCC), and austenite has a face-centered cubic structure. As can be seen from the figure, the diffraction peaks of the reverse-phase transformation austenite first increase and then decrease with increasing annealing time, reaching their highest values ​​at 800℃-5min / 640℃-15min. Figure 19cThe austenite content trends of the two groups of experimental steels are consistent. When the holding time is 5 min, the austenite content of the two groups of experimental steels is 18.9% and 20.5% respectively. With the extension of holding time, the austenite content increases and reaches the highest at 15 min, which is 24.7% and 25.9% respectively. When the annealing time is extended to 25 min, the austenite content decreases, which is 21% and 20.8% respectively for the two groups of experimental steels. This is because the austenite grains grow with the increase of holding time, and the stability of austenite decreases, so a part of austenite is transformed into martensite during cooling. This also confirms the evolution law of austenite in the microstructure of the experimental steels. By comparing the austenite volume fraction of the two groups of experimental steels under different ART annealing time, it can be found that the addition of rare earth elements in the steel can improve the stability of austenite, so that more residual austenite structure can be obtained at room temperature.

[0257] The mechanical properties of the experimental steels treated by ART process with different annealing time are tested, and the results are shown in Figure 20a- Figure 20f

[0258] As can be seen from Figure 20a- Figure 20f , the yield strength of the two groups of experimental steels gradually decreases with the extension of holding time. This is because the yield strength is mainly affected by the grain size. The longer the holding time, the larger the grain size, and the smaller the yield strength. When the holding time is 5 min, the yield strength of the experimental steels without rare earth and with rare earth is the highest, which is 993 MPa and 1028 MPa respectively. The tensile strength gradually increases with the extension of holding time. In the case of short holding time, the tensile strength is mainly affected by TRIP effect. The more the austenite content in the steel, the longer the TRIP effect lasts, and the higher the tensile strength. When the holding time is 15 min, the austenite content is the highest, but the tensile strength is not the highest. This is because with the further extension of holding time, the austenite grains grow and become unstable, and a part of them is transformed into martensite during cooling. Under the joint action of martensite strengthening and TRIP effect, the tensile strength increases. When the holding time is 25 min, the tensile strength is the highest, which is 1289 MPa and 1189 MPa respectively. The elongation and the product of strength and plasticity both increase first and then decrease with the extension of annealing time, and reach the highest at 15 min. The elongation of the two groups of experimental steels is 32.3% and 33.2% respectively, and the product of strength and plasticity is 33.1 GPa·% and 34.8 GPa·% respectively. The elongation is related to the austenite content. When the holding time is 15 min, the austenite content is the highest, and the TRIP effect lasts for a long time, so that the austenite transforms into martensite during deformation, the necking is delayed, and thus the elongation increases.

[0259] Figure 21a- Figure 21f The fracture morphologies of the two groups of experimental steels after quenching at 800℃ for 5 min and different annealing time.

[0260] As can be seen from​Figure 20a- Figure 20f It can be seen that the elongation of the material is different, and the fracture behavior is also different, and the fracture morphology can be observed to determine whether the material is brittle or ductile. Figure 21a 、 Figure 21d The tensile fracture morphology of the annealing for 5 min can be seen that there are fewer dimples, and there are parting fractures. When the annealing time is extended to 15 min, as shown in Figure 21b 、 Figure 21e , it is basically a ductile fracture, and the dimples are more densely and uniformly distributed, and the elongation of the material is the highest, i.e. the plasticity is the best. When the holding time is 25 min, as shown in Figure 21c 、 Figure 21f , the fracture morphology of the sample is also basically a ductile fracture, but the dimples are not uniformly distributed, and the plasticity is lower than that of 15 min. By comparing the fracture morphology of the two groups of experimental steels, it is found that the dimples of the experimental steel containing rare earth are smaller, more uniform and dense, and it is inferred that the addition of rare earth improves the plasticity of the material, which is consistent with the change rule of the microstructure and mechanical properties of the above-mentioned material.

[0261] By annealing two groups of experimental steels containing 48 ppm of rare earth Ce and not containing rare earth by ART, the technician studies the evolution rule of annealing temperature and annealing time on the microstructure, mechanical properties and forming properties, and further studies the effect of rare earth micro-alloying on the microstructure and properties of ultra-high strength automobile steel, and the summary is as follows:

[0262] (1) The addition of trace rare earth elements increases the Ac3 temperature of the experimental steel, which delays the nucleation during austenitization, resulting in finer original austenite grains of rare earth automobile steel.

[0263] (2) The strength-plasticity product of the hot-rolled plate and the cold-rolled plate without rare earth is 13.7 GPa·% and 12.2 GPa·%, respectively, and the strength-plasticity product of the hot-rolled plate and the cold-rolled plate containing rare earth is 16.2 GPa·% and 13.9 GPa·%, respectively. The addition of rare earth elements can improve the comprehensive mechanical properties.

[0264] (3) After ART annealing, the microstructure of the experimental steels is composed of ferrite, austenite and newly formed martensite. The grain size distribution of the experimental steels with rare earth elements is more uniform, and the comprehensive mechanical properties are better. With the increase of annealing temperature, the tensile strength gradually increases, and the elongation and the product of strength and elongation first increase and then decrease. The comprehensive mechanical properties and forming properties are the best when annealed at 800℃-5min / 640℃-15min. The residual austenite content of the experimental steel without rare earth and with rare earth is 24.7% and 25.7% respectively, the tensile strength is 1024MPa and 1048MPa respectively, the elongation is 32.3% and 33.2% respectively, and the product of strength and elongation is 33.1GPa·% and 34.8GPa·% respectively. The influence of different ART annealing time on the experimental steel is consistent with the annealing temperature.

[0265] 3. Effect of Ce on microstructure and properties of ultra-high strength automobile steel under two-phase zone direct annealing process

[0266] 3.1 Effect of two-phase zone direct annealing process on microstructure and properties of rare earth microalloyed ultra-high strength automobile steel

[0267] 3.1.1 Effect of annealing temperature on microstructure and properties

[0268] In order to compare with ART process, the annealing temperature of two-phase zone direct annealing is the same as that of ART process, and the temperature is 620℃, 640℃ and 660℃ for 15min and then air-cooled to room temperature. Figure 22a- Figure 22f The microstructure of the experimental steel annealed at different temperatures in two-phase zone is shown in the following figures.

[0269] When the annealing temperature is 620℃, as shown in Figure 22a , Figure 22d , the microstructure is composed of ferrite and austenite. The convex bright white structure is austenite, and the concave dark gray structure is ferrite. Most of the structure is lath-shaped, and a small part is equiaxed. This is because the martensite structure is transformed into lath-shaped ferrite and austenite at 620℃, but the microstructure cannot be completely equiaxed for 15min. When the annealing temperature is increased to 640℃, as shown in Figure 22b , Figure 22e , the microstructure is still composed of austenite and ferrite. Compared with 620℃, the microstructure is coarsened, the equiaxed degree is increased, and the lath-shaped austenite begins to fuse together, and the austenite content is increased. When the annealing temperature is 660℃, as shown in Figure 22c , Figure 22fAs shown, the microstructure at this time is a multi-phase structure composed of ferrite, martensite and residual austenite, because the annealing temperature is increased, the grain grows, resulting in the decrease of C and Mn element content inside austenite grain, and the stability of austenite also decreases, a part of austenite is transformed into martensite during the cooling process, and the lath-shaped austenite is obviously coarsened and tends to develop into equiaxed. By observing the microstructure of the two groups of experimental steels, it is not difficult to find that the protruding structure in the experimental steel with rare earth is obviously more than that without rare earth, which shows that the addition of rare earth plays a certain role in increasing the content of austenite.

[0270] Figure 23a- Figure 23c TEM images of the two groups of experimental steels after annealing at 640℃, Figure 23a 、 Figure 23b The TEM images of the experimental steel without rare earth and with rare earth respectively, as can be seen from the figure, the microstructure after annealing is composed of ferrite and austenite, and there are dislocations inside the ferrite, as shown in Figure 23c , these dislocations are formed by the phase transformation from face-centered cubic austenite to body-centered cubic ferrite at high temperature, and the existence of dislocations can improve the strength of the material. Figure 24a 、 Figure 24c Austenite TEM images of the experimental steel without rare earth and with rare earth respectively, Figure 24b 、 Figure 24d Austenite diffraction spots of the experimental steel without rare earth and with rare earth respectively, the austenite grain size of the experimental steel without rare earth is 0.5μm, and the austenite grain size of the experimental steel with rare earth is 0.3μm, it is found by comparison that the addition of rare earth can refine the grain.

[0271] The volume fraction and stability of austenite have an important influence on the mechanical properties of the experimental steel, the strength of TRIP effect depends on the volume fraction of austenite, and the duration of TRIP effect is related to the stability of austenite, therefore the volume fraction and stability of austenite directly determine the plasticity of the steel. Figures 25a-25c XRD patterns and austenite volume fraction diagrams of the two groups of experimental steels after direct annealing at different temperatures in the two-phase region, the γ peak represents face-centered cubic structure, and the α peak represents body-centered cubic structure, the volume fraction of austenite can be calculated by the peak intensity of inverse phase austenite and ferrite. As shown in Figure 25a 、 Figure 25b , the austenite diffraction peak first increases and then decreases, and reaches the highest at 640℃, indicating that the content of residual austenite obtained by annealing at 640℃ is the highest, the volume fraction of austenite of the experimental steel without rare earth and with rare earth respectively reaches 26.7% and 25.8%, when the temperature continues to rise to 660℃, the γ peak value decreases, and the α peak value increases obviously, because the temperature is increased, a part of austenite is transformed into martensite when cooling to room temperature, and martensite is also body-centered cubic structure like ferrite, therefore the α peak value representing body-centered cubic structure increases obviously, and the austenite diffraction peak decreases.Figure 25c As shown in the figure, the austenite volume fraction of the experimental steel containing rare earth is higher than that of the experimental steel without rare earth, which indicates that the mechanical properties are enhanced after adding rare earth, which is consistent with the change rule of microstructure.

[0272] Figures 26a-26f The stress-strain curves and the mechanical property change curves of the two groups of experimental steels under different critical annealing temperatures are shown in the figure.

[0273] As shown in the figure, the comprehensive index of each mechanical property of the sample is the best when the annealing temperature is 640℃, at which the yield strength of the experimental steel without rare earth and containing rare earth is 1062MPa and 1125MPa respectively, the tensile strength is 1230MPa and 1265MPa respectively, the elongation is 32.3% and 33.2% respectively, and the strength-plasticity product is 38GPa·% and 40.2GPa·% respectively. The yield strength of the experimental steel gradually decreases with the increase of the annealing temperature. The first reason is that the lath-shaped structure in the steel tends to be equiaxed with the increase of the temperature, and the back stress strengthening is formed by the back stress generated by the dislocation accumulation of ferrite during deformation. With the increase of the annealing temperature, the volume fraction of ferrite decreases, and studies have shown that the dislocation of equiaxed structure proliferates along the grain boundary at the beginning of deformation, and the dislocation density generated inside the grain is lower than that at the grain boundary, so the dislocation density of lath-shaped ferrite is higher than that of equiaxed ferrite, therefore the yield strength gradually decreases with the increase of the annealing temperature. The second reason is that the grain size gradually grows with the increase of the annealing temperature, and according to formula 3-1, the larger the grain size, the lower the yield strength. The tensile strength of the experimental steel gradually increases with the increase of the annealing temperature. When the annealing temperature is 620℃, the grain size of the sample is the smallest, the austenite content is the lowest, and the TRIP effect is the weakest. When the annealing temperature rises to 640℃, the grain size increases, the austenite content increases, and the degree of martensite transformation during tensile process is larger than that at 620℃, so the tensile strength increases due to the martensite transformation strengthening. When the annealing temperature rises to 660℃, the austenite content decreases and the degree of martensite transformation decreases, but a part of austenite is transformed into martensite during cooling, so the tensile strength increases.

[0274]

[0275] In the formula: σS is the yield strength; σ0 is the friction force acting on the dislocation; d is the average grain diameter.

[0276] As Figure 26eAs shown, the elongation of the experimental steel first increased and then decreased with increasing temperature, reaching its highest value at 640℃. When the annealing temperature was 620℃, the austenite stability was very high, and no martensitic transformation occurred during deformation, resulting in a weakened TRIP effect. When the annealing temperature increased to 640℃, not only lath austenite but also some blocky and equiaxed austenite were present. Since the stability of blocky and equiaxed austenite is lower than that of lath austenite, the TRIP effect is enhanced. Moreover, the austenite volume fraction is highest at 640℃, resulting in the highest elongation. When the annealing temperature increased to 660℃, the residual austenite content decreased, becoming comparable to that at 620℃, but the austenite stability decreased. During deformation, almost all austenite transformed into martensite, so the elongation at 660℃ was higher than that at 620℃. The trend of the strength-ductility product change was consistent with the elongation, reaching its highest value at 640℃.

[0277] like Figures 27a-27f The image shows the fracture morphology of two groups of experimental steels after direct annealing in the two-phase region at different annealing temperatures.

[0278] The fracture morphology can be characterized by cleavage and dimples. When the annealing temperature is 620℃, as shown in the figure... Figure 27a , Figure 27d As shown, the fracture morphology at this point is basically a river-like cleavage fracture, combined with... Figures 26a-26f The mechanical property change curve shows that the plasticity is worst at 620℃. When the temperature rises to 640℃, as... Figure 27b , Figure 27e As shown, at this point, the fracture surface is mainly composed of dimples, which are evenly and densely distributed, indicating the best plasticity. When the temperature continues to rise to 660℃, as... Figure 27c , Figure 27f As shown, the fracture surface consists of dimples and river-like cleavage planes, therefore the plasticity at 660℃ is between 620℃ and 640℃. By comparing the fracture morphology of the two groups of experimental steels, it can be found that the dimple distribution of the rare earth-containing experimental steel is more uniform, thus indicating that the change in elongation is a manifestation of the distribution and uniformity of dimples.

[0279] 3.1.2 Effect of annealing time on microstructure and properties

[0280] Figures 28a-28f The microstructures of two groups of experimental steels were obtained after direct annealing in the two-phase region at 640℃ for different annealing times: 5 min, 15 min, and 25 min. The transformation from martensite to austenite is very rapid due to the high-density nucleation sites of austenite in the lath martensite matrix, and because the chemical potential of carbon in martensite is higher than that in austenite, carbon atoms diffuse continuously from martensite to austenite. When the annealing time is 5 min, as... Figure 28a , Figure 28dAs shown, the microstructure is a multiphase structure of lath ferrite, austenite, and martensite. Due to the short annealing time and low degree of carbon atom diffusion, the martensite in the experimental steel did not completely transform into austenite. When the annealing time is extended to 15 minutes, as... Figure 28b , Figure 28e As shown, the raised structure consists of austenite and martensite. At this point, austenite nucleates between the martensite laths, and the martensite changes from a lath-like shape to a blocky shape. (See diagram.) Figure 28c , Figure 28f As shown, with further extension of annealing time, both martensite and austenite grains gradually become equiaxed.

[0281] The austenite content of two groups of experimental steel samples annealed at 640℃ for different times was analyzed by X-ray diffraction (XRD) to clarify the influence of annealing time on the microstructure of the experimental steel. The results are as follows: Figures 29a-29c As shown in the figure, the γ peak represents the face-centered cubic structure, and the α peak represents the body-centered cubic structure. When the annealing time is 5 min, the austenite content of the experimental steels without rare earth elements and those containing rare earth elements is 20.3% and 21.9%, respectively. Critical annealing can cause some martensite in the microstructure to undergo reverse phase transformation to obtain reverse-transformed austenite. However, due to the short holding time, the obtained reverse-transformed austenite has high stability, and during the subsequent tensile process, the amount of martensite phase change is small, and the TRIP effect is not obvious. With the extension of the holding time, the diffraction peaks of the reverse-transformed austenite first increase and then decrease, indicating that the austenite content first increases and then decreases, reaching its highest point at 15 min of holding time, with the content of the experimental steels without rare earth elements and those containing rare earth elements being 25.9% and 26.4%, respectively. When the annealing time is extended to 25 min, the α diffraction peak is significantly enhanced. This is because the holding time is too long, which leads to a decrease in austenite stability, and some austenite transforms into martensite during the cooling process. By comparing the austenite content of the two groups of experimental steels, it was found that the austenite content of the experimental steel containing rare earth elements was higher than that of the experimental steel without rare earth elements. This indicates that the addition of rare earth elements is beneficial to enhancing the stability of retained austenite, allowing the experimental steel to retain more retained austenite structure at room temperature and improving the mechanical properties of the steel.

[0282] like Figures 30a-30f The stress-strain curves and mechanical property change curves of two groups of experimental steels after annealing at 640℃ for different times are shown.

[0283] like Figure 30cAs shown, the yield strength of both groups of experimental steels gradually decreased with the extension of annealing time. This is because as the annealing time increases, the grains of the experimental steel grow continuously, and the yield strength is mainly related to the grain size; the larger the grain, the lower the yield strength. The tensile strength gradually increased with the extension of annealing time. The tensile strength is mainly affected by the TRIP effect and martensitic strengthening. When the annealing time is 5 min, the austenite content is low and its stability is high, and the TRIP effect is weak, resulting in the lowest tensile strength. When the annealing time is 15 min, the austenite content is highest, and the TRIP effect is significant, leading to an increase in tensile strength. If the annealing time is further extended to 25 min, the austenite becomes unstable and partially transforms into martensite during cooling to room temperature. Therefore, the combined effect of martensitic strengthening and the TRIP effect results in the highest tensile strength. At this point, the tensile strengths of the experimental steels without rare earth elements and those containing rare earth elements are 1155 MPa and 1247 MPa, respectively. The elongation first increased and then decreased with the extension of annealing time. The elongation was significantly affected by the TRIP effect. When the TRIP effect occurs, unstable austenite transforms into martensite, resulting in delayed necking. The higher the austenite content, the longer the TRIP effect lasts, and the higher the elongation. The change pattern of the strength-ductility product was consistent with the elongation, reaching its highest point at 15 min. The elongation of the experimental steels without rare earth elements and those containing rare earth elements were 32.3% and 33.2%, respectively, and the strength-ductility products were 38 GPa·% and 40.2 GPa·%, respectively. Comparing the mechanical properties of the two groups of experimental steels, it was found that the experimental steel containing rare earth elements had better mechanical properties.

[0284] like Figures 31a-31f The fracture morphology is shown for two-phase direct annealing at different times. Fracture is the endpoint of continuous deformation, therefore the fracture morphology can reflect the characteristics of the material deformation process. During tensile testing, ferrite is continuously elongated, causing grain to twist. Dislocations within the grains continuously move and entangle, leading to stress concentration. When stress acts on the retained austenite, the austenite undergoes a martensitic transformation, i.e., the TRIP effect. After the TRIP effect ends, continued deformation causes martensite to be pulled out of the ferrite, forming dimples. The size, distribution, and uniformity of the dimples can reflect the plasticity of the material. When the annealing time is 5 minutes, for example... Figure 31a , Figure 31d As shown, the fracture morphology is a mixture of cleavage and dimples. When the annealing time is increased to 15 minutes, as... Figure 31b , Figure 31e As shown, the fracture morphology at this point is basically dimples, and they are evenly and densely distributed, indicating that the material has good plasticity under this heat treatment process. When the annealing time is increased to 25 minutes, as... Figure 31c , Figure 31fAs shown in the figure, the fracture morphology at this time is a mixture of dimples and cleavage, but the dimples are more distributed than those at 5 min. Through the analysis of the fracture morphology, it can be inferred that the material has the best plasticity when the annealing time is 15 min, which is consistent with the mechanical property analysis.

[0285] 3.2 Comparative analysis of two annealing processes

[0286] Through the study of two heat treatment processes for two groups of experimental steels, the best heat treatment process parameters are found: the best heat treatment process parameters for two-phase zone direct annealing are 640℃-15min, and the best heat treatment process parameters for ART process are 800℃-5min+640℃-15min. By comparing the microstructure and mechanical properties under the two best heat treatment process parameters, and testing the texture, it is analyzed which heat treatment process is more beneficial to forming.

[0287] 3.2.1 Microstructure analysis

[0288] Figures 32a-32d and Figures 33a-33d are SEM and TEM images of two groups of experimental steels under the best heat treatment conditions, as shown in Figures 32a-32d , the dark structure in the TEM image is austenite, and the light-colored one is ferrite structure. The experimental steel after ART annealing has more carbides and poor uniformity of structure. From the TEM image, it can be seen that the austenite structure after ART annealing is mostly lath-shaped, while the austenite structure after two-phase zone annealing is mostly block-shaped.

[0289] Figures 34a-34d are TEM images under the two best heat treatment processes. From the figure, it can be found that there are dislocations in the austenite and ferrite under the two processes, but the dislocation density in the ferrite after two-phase zone direct annealing is higher. This is because in two-phase zone annealing, the crystal is divided into two phases, so the temperature gradient will be larger, which may cause some parts of the crystal to cool too fast, thereby forming more dislocations. In ART annealing, the temperature gradient is relatively small, so it is not easy to produce such rapid cooling conditions, thereby reducing the generation of dislocations. Figure 5 is a graph of the relationship between dislocation density and strength. From the figure, it can be seen that the dislocation of the alloy after heat treatment is proportional to the strength. Since there are more dislocations in the ferrite of the experimental steel after two-phase zone direct annealing, its strength is higher than that of the experimental steel after ART annealing.

[0290] 3.2.2 Mechanical property comparison

[0291] Figures 36a-36bFor the comparison of mechanical properties of two annealing processes, it can be seen from the figure that the yield strength of the material annealed in the two-phase region is higher than that of the ART process, because the grain size after two-phase region annealing is smaller than that of ART process, and the yield strength is mainly affected by the grain size, the larger the grain size, the smaller the yield strength. The tensile strength of the material after two-phase region annealing is also higher than that of the ART process, because the ART process has a quenching process before annealing, and the microstructure after quenching is lath-shaped martensite, and the stability of the residual austenite obtained after reverse phase transformation annealing is higher than that of the residual austenite obtained after two-phase region annealing, which leads to lower degree of martensitic transformation during tensile process and weaker TRIP effect. Another reason is that there are a large number of dislocation pile-ups in the ferrite structure after two-phase region annealing, which will cause local stress concentration, making it difficult for the material to deform under external force, so the experimental steel annealed in the two-phase region has higher tensile strength. By comparing the elongation of the two processes, it is found that the elongation of the experimental steel after ART annealing is higher than that of the two-phase region annealing, because the bulk austenite is a large block structure based on cubic crystal system, while the lath-shaped austenite is a metastable structure composed of fine fibrous grains. Because the lattice structure of bulk austenite is relatively large, it is more prone to fracture; while the lath-shaped austenite has a fine structure due to its slender lattice structure, thus has higher elongation. Although the elongation of the experimental steel after ART annealing is high, the product of strength and plasticity is lower than that of the experimental steel after two-phase region annealing, so two-phase region annealing can replace ART annealing, that is, it can save production cost and obtain higher comprehensive mechanical properties.

[0292] 3.2.3 Forming property analysis

[0293] The forming property of automobile steel is related to the plastic strain ratio r value, the larger the r value, the better the forming property of the steel sheet, and the r value is closely related to the texture density of the material. By testing the texture density and content of the experimental steel under different heat treatment processes, the influence of heat treatment process on the forming property can be judged. The texture density and content of the experimental steel can be analyzed by orientation distribution function (ODF graph), and the orientation line density in the ODF graph represents the orientation density. Under normal circumstances, only the most important orientation distribution change needs to be analyzed, and the grain orientation of the cold rolled sheet gradually converges to the α and γ orientation lines during rolling, therefore, this study mainly analyzes the texture type, texture density and texture content on the two orientation lines.

[0294] Figures 37a-37d For the ODF section graph of φ2 = 45° of the two groups of experimental steel under different heat treatment processes, it can be seen from the figure that the two groups of experimental steel after different heat treatment processes all exist {001} <110>, {011} <110>, {111} <110> and {111} <112> textures, Figure 37a ,Figure 37c For the experimental steels annealed by ART, we found that both groups of experimental steels had {111} on the γ orientation line. <110> Strong peaks appear near the texture, and the texture distribution of the experimental steel containing rare earth elements is more uniform, with a maximum texture density of 2.49. The texture density of the experimental steel annealed directly in the two-phase region is greater than that of the ART process, with a maximum texture density of 4.02.

[0295] Tables 10 and 11 show the texture content of the two groups of experimental steels under different heat treatment conditions for each orientation. It can be seen from the tables that the texture content of the two groups of experimental steels does not change significantly in each orientation. Figure 36a , Figure 36b Analysis shows that adding trace amounts of rare earth elements to automotive steel can improve its formability, but the improvement is not significant. Further analysis is needed by calculating the favorable texture {111} and unfavorable texture {100}. By comparing the texture fractions of the experimental steels under two processes, it was found that the texture of {001} directly annealed in the two-phase region... <110> {112} <110> {223} <110> {111} <110> {111} <112> The volume fraction of the texture is higher than that of the ART process, therefore the experimental steel annealed directly in the two-phase region has better formability.

[0296] Table 10 Texture content of different orientations under ART annealing process

[0297]

[0298] Table 11 Texture content of different orientations under two-phase direct annealing process

[0299]

[0300] The {111}, {110}, and {100} texture contents of the two groups of experimental steels under different heat treatment conditions were calculated using ResMat-TexTools software. Figure 38a , Figure 38b As shown, the {111} texture can improve the formability of steel sheets, while the {100} texture is detrimental to the formability of steel sheets, and the {110} texture is between the {111} and {100} textures.

[0301] like Figure 38a As shown, for the ART annealing process, the {111} texture content of the experimental steels without rare earth elements and those containing rare earth elements were 12.3% and 13.4%, respectively; the {110} texture content was 19.5% and 20.3%, respectively; and the content of the unfavorable texture {100} was 13.8% and 11.6%, respectively. Figure 38bAs shown, for the two-phase region direct annealing process, the {111} texture content of the experimental steel without rare earth and with rare earth is 17.1% and 17.4% respectively, the {110} texture content is 24.5% and 25.3% respectively, and the content of the adverse texture {100} is 11.2% and 10.1% respectively. By comparison, it is found that the {111} and {110} texture content of the experimental steel containing rare earth is higher than that of the experimental steel without rare earth, and the adverse texture {100} of the experimental steel without rare earth is higher than that of the experimental steel containing rare earth. By comparing the texture content under two heat treatment processes, it is found that the {111} and {110} texture content after two-phase region direct annealing is higher, and the content of the adverse texture {100} is lower, so it can be seen that the forming property of the experimental steel after two-phase region direct annealing is better than that of ART annealing process.

[0302] This part mainly carries out two-phase region direct annealing on two groups of experimental steels containing 48ppm rare earth Ce and without rare earth, studies the evolution law of annealing temperature and annealing time on microstructure, mechanical properties and forming properties, explores the influence of rare earth micro-alloying on the microstructure and properties of ultra-high strength automobile steel, and summarizes as follows:

[0303] (1) After two-phase region direct annealing, the microstructure of the experimental steel is composed of ferrite, austenite and new martensite, and the grain size distribution of the experimental steel after adding rare earth element is more uniform, and the comprehensive mechanical properties are better. With the increase of annealing temperature, the tensile strength gradually increases, the elongation and the product of strength and plasticity first increase and then decrease, and the comprehensive mechanical properties and forming properties are the best when annealed at 640℃ for 15min, the residual austenite content of the experimental steel without rare earth and with rare earth is 25.8% and 26.7% respectively, the tensile strength is 1230MPa and 1265MPa respectively, the elongation is 30.9% and 31.8% respectively, and the product of strength and plasticity reaches 38GPa·% and 40.2GPa·% respectively. The influence law of different annealing time on the experimental steel is consistent with that of annealing temperature.

[0304] (2) The trace rare earth element has an influence on the texture content, and the texture distribution and density level of the experimental steel containing rare earth element and without rare earth element are similar, but the favorable texture {111}, {110} volume fraction of the experimental steel containing rare earth element is higher than that without rare earth element, and the unfavorable texture {100} volume fraction is lower, so it can be seen that the trace rare earth element not only affects the mechanical properties of the experimental steel, but also significantly improves the formability of the experimental steel after adding rare earth element.

[0305] (3) By comparing the texture content and texture density of the ultra-high strength automobile steel under two heat treatment processes, it is found that the texture content after two-phase region direct annealing is higher, which reveals that the comprehensive mechanical properties and forming properties of the experimental steel after two-phase region direct annealing are better.

[0306] 4. Deformation process research of rare earth micro-alloyed ultra-high strength automobile steel

[0307] This part will study the deformation process of the experimental steel, including deformation temperature, deformation degree and deformation speed. But the experimental steel studied in this part is cold-rolled sheet, and the effect of deformation speed on its microstructure and properties is not great, so this part will focus on the effect of deformation degree on the microstructure and properties of rare earth micro-alloyed ultra-high strength automobile steel. By controlling the deformation degree of the steel sheet, the mechanical properties of the steel sheet can be changed, including strength, hardness and ductility, etc. Controlling the deformation degree of the steel sheet can also improve production efficiency and reduce manufacturing cost. This part will study the effect of deformation degree on the microstructure and properties of 48ppmRE experimental steel with deformation degree of 73%, 50% and 33% after two-phase region direct annealing (640℃-15min).

[0308] 4.1 Effect of deformation degree on microstructure and texture

[0309] Figures 39a-39c The microstructure photos of the experimental steel after annealing with different deformation degrees. From the figure, it can be seen that the microstructure of the experimental steel with deformation degree of 73%, 66% and 33% after annealing at 640℃ for 15min is dual-phase microstructure of austenite and ferrite. With the increase of deformation degree, the grain of the steel sheet is gradually refined. This is because during rolling, the steel sheet will undergo plastic deformation, and the crystal structure will also change. The original regular grains are elongated and produce a large number of dislocations, which in turn trigger grain recrystallization, resulting in grain refinement. Grain refinement will increase the strength and hardness of the steel sheet, but the plasticity will decrease accordingly. It can be inferred that the comprehensive mechanical properties change little.

[0310] The forming property of steel sheet is related to the r-value, the greater the r-value, the better the forming property of the material, the less likely to wrinkle in the stamping process, which can improve the utilization of the material, and also can avoid cracks in stamping. The size of r-value is affected by many factors, such as impurity element content, processing technology, heat treatment process and grain size, etc., but the most important influencing factor is the texture of the steel sheet. The reason why IF steel has the intended deep drawing property is that it has a strong {111} plane texture, also called γ silk texture. In the {111} plane texture, the grains are arranged in the same direction, and this type of crystal direction has high density and uniformity, which leads to excellent performance in mechanical properties. For the stamping process, the most important is the ductility and work hardening of the material, the grain orientation of {111} plane texture can enhance these two mechanical properties, so that the more brittle stamping material can be successfully formed. Some documents show that {111} plane texture can increase the r-value of the material, while {100} plane texture can reduce the r-value, because the {111} crystal plane is the main slip plane, and the <100> direction is the main slip direction, this slip system is parallel to the plate surface, and has strong resistance to thickness reduction when the steel sheet is deformed, so the deep drawing property is good. For example Figures 40a-40c The texture of cold rolled sheet is generally concentrated in the α and γ direction lines, in order to clearly observe the texture in the two direction lines, the φ2=45° ODF texture diagram of different deformation degrees is drawn out, as shown in Figures 41a-41c It can be seen from the figure that the experimental steel of different deformation degrees exists in {001}<110>, {112}<110>, {223}<110>, {111}<110> and {111}<112> texture.

[0311] As shown in Table 12, the texture volume fraction of <110> and {111} of the experimental steel of different deformation degrees is calculated by software, it can be found that with the increase of deformation degree, the volume fraction of {001}<110>, {112}<110>, {223}<110>, {111}<110>, {111}<112> texture gradually increases. Therefore, from the volume fraction of each orientation texture, the experimental steel with a deformation degree of 73% has the best forming property.

[0312] Table 12 Texture volume fraction of <110> and {111} of experimental steel of different deformation degrees

[0313]

[0314] The {111}, {110} and {100} texture content of the experimental steel of different deformation degrees after two-phase region direct annealing is calculated by ResMat-TexTools texture analysis software, as shown inFigure 42 As shown in the figure, with the increase of deformation degree, the content of favorable textures {111} and {110} gradually increases, and the content of unfavorable texture {100} gradually decreases. The r value is the largest when the deformation degree is 73%, i.e. the forming performance is the best.

[0315] 4.2 Influence of deformation degree on performance

[0316] Figure 43 The stress-strain curves of the 48 ppm RE experimental steel with different deformation degrees (direct annealing in two-phase region at 640℃ for 15 min) are shown in the figure. As can be seen from the figure, the mechanical properties of the experimental steel with three different deformation degrees are similar. As shown in Table 13, the calculated product of strength and ductility of the three experimental steels is 38.6 GPa·%, 40.2 GPa·% and 39.7 GPa·% respectively. Although the mechanical properties are not much different, the r value gradually increases with the increase of deformation degree. Therefore, when the deformation degree is 73%, the forming performance is the best, which is consistent with the change of texture described above.

[0317] Table 13 Mechanical properties of experimental steel with different deformation degrees after heat treatment

[0318]

[0319] This part studies the influence of different deformation degrees of cold-rolled plates on the microstructure, texture, mechanical properties and forming performance of automobile steel, and the summary is as follows:

[0320] (1) With the increase of deformation degree, the grains of the experimental steel are gradually refined, the tensile strength gradually increases, the elongation gradually decreases, but the product of strength and ductility changes little, and the mechanical properties of automobile steel with different deformation degrees are similar.

[0321] (2) With the increase of deformation degree, the favorable textures {111} and {110} of the experimental steel continuously increase, the unfavorable texture {100} gradually decreases, and the r value also gradually increases with the increase of deformation degree. The larger the r value, the better the forming performance of the material, which shows that the increase of deformation degree can make the comprehensive performance of the material better. Therefore, the experimental steel with a deformation degree of 73% has better forming performance after 640℃-15min.

[0322] 5. Conclusion

[0323] The present embodiment designs the ultra-high strength automobile steel containing 48 ppm Ce and without rare earth in the basis of 0.12C-5Mn system manganese steel, and studies the ART annealing and two-phase zone direct annealing heat treatment process, so as to realize the target of high strength and high plasticity. The change rule of annealing time and annealing temperature on the microstructure, mechanical properties and texture of the ultra-high strength automobile steel is studied by means of SEM, TEM and XRD technology, and the influence of rare earth micro-alloying on the microstructure and properties of the ultra-high strength automobile steel is explored. On this basis, the influence of deformation degree on the microstructure and properties of the ultra-high strength automobile steel is analyzed, and the most favorable deformation degree for forming is determined. The conclusions are as follows:

[0324] (1) The addition of rare earth Ce has a refining effect on the original austenite grains of the ultra-high strength automobile steel, and the average size of the original austenite grains of the experimental steel containing rare earth and without rare earth is 4.42 μm and 4.91 μm respectively, because the rare earth elements can form inclusions in the grain to hinder the grain growth, thereby achieving the effect of refining the grains.

[0325] (2) The product of strength and plasticity of the hot-rolled plate and the cold-rolled plate without rare earth is 13.7 GPa·% and 12.2 GPa·% respectively, and that of the experimental steel containing rare earth is 16.2 GPa·% and 13.9 GPa·% respectively, and the addition of rare earth elements improves the mechanical properties of the hot-rolled plate and the cold-rolled plate.

[0326] (3) In the ART annealing process, the microstructure of the experimental steel after quenching is lath-shaped martensite, and the microstructure of the experimental steel after ART annealing is mainly ferrite, residual austenite and newly generated martensite. With the increase of annealing temperature and time, the grains of the two groups of experimental steel gradually grow, the yield strength gradually decreases, the tensile strength gradually increases, and the elongation and the product of strength and plasticity first increase and then decrease. When the annealing temperature is low, a large amount of carbide is produced in the microstructure, and with the increase of temperature, the carbide gradually decreases and disappears, and when the temperature is high, new martensite appears in the microstructure, which is transformed from unstable austenite during cooling. When annealing at 800℃-5min / 640℃-15min, the comprehensive mechanical properties and forming properties are the best, the residual austenite content of the experimental steel without rare earth and containing rare earth is 24.7% and 25.7% respectively, the tensile strength is 1024 MPa and 1048 MPa respectively, the elongation is 32.3% and 33.2% respectively, the product of strength and plasticity reaches 33.1 GPa·% and 34.8 GPa·% respectively, the volume fraction of {111} texture is 12.3% and 13.4% respectively, the volume fraction of {110} texture is 19.5% and 20.3% respectively, and the volume fraction of the adverse texture {100} is 13.8% and 11.6% respectively, and the addition of rare earth improves the mechanical properties and forming properties of the automobile steel.

[0327] (4) In the two-phase region direct annealing process, with the increase of annealing temperature and time, the grain of the experimental steel gradually equiaxed, the tensile strength gradually increased, the content of residual austenite first increased and then decreased, the mechanical properties related to the content of residual austenite, the change rule of elongation and product of strength and elongation was consistent with the change rule of residual austenite. The comprehensive mechanical properties and forming properties were the best when annealing at 640℃-15min, the residual austenite content of the steel without rare earth and the steel containing rare earth was 25.8% and 26.7% respectively, the tensile strength was 1230MPa and 1265MPa respectively, the elongation was 30.9% and 31.8% respectively, the product of strength and elongation reached 38GPa·% and 40.2GPa·% respectively, the volume fraction of {111} texture was 17.1% and 17.4% respectively, the volume fraction of {110} texture was 24.5% and 25.3% respectively, the volume fraction of adverse texture {100} was 11.2% and 10.1% respectively, the addition of rare earth improved the mechanical properties and forming properties of the automobile steel.

[0328] (5) Under the optimum conditions of two kinds of heat treatment, the microstructure of ART annealing was less uniform than that of two-phase region direct annealing; the elongation of ART annealing was higher, which was because the austenite microstructure of ART annealing was strip-shaped and the austenite microstructure of two-phase region direct annealing was block-shaped, but the yield strength, tensile strength and product of strength and elongation were lower than those of two-phase region direct annealing, the product of strength and elongation of the experimental steel containing rare earth and the experimental steel without rare earth after ART annealing was 33.1GPa·% and 34.8GPa·% respectively, the product of strength and elongation of the experimental steel containing rare earth and the experimental steel without rare earth after two-phase region direct annealing was 38GPa·% and 40.2GPa·% respectively; the volume fraction of favorable texture of the experimental steel after ART annealing was lower than that of two-phase region direct annealing, while the volume fraction of adverse texture was higher than that of two-phase region direct annealing, which indicated that the two-phase region direct annealing process was more beneficial to the forming of ultra-high strength automobile steel. It can be seen that the rare earth micro-alloyed ultra-high strength automobile steel after two-phase region direct annealing can achieve good mechanical properties and forming properties, and can also save cost and avoid resource waste.

[0329] (6) With the increase of the deformation degree, the grain of the rare earth micro-alloyed experimental steel is gradually refined, the tensile strength gradually increases, the elongation gradually decreases, but the product of strength and ductility changes little, and the mechanical properties of the automobile steels with different deformation degrees are similar; the influence of the deformation degree on the texture and formability is analyzed. The results show that with the increase of the deformation degree, the beneficial textures {111} and {110} of the experimental steel continuously increase, the harmful texture {100} gradually decreases, and the r value gradually increases with the increase of the deformation degree, the larger the r value, the better the formability of the material, which shows that the increase of the deformation degree can make the comprehensive performance of the material better. Therefore, the experimental steel with a deformation degree of 73% has the best formability after direct annealing in the two-phase region at 640℃ for 15min, at this time, the {111} texture content is 17.8%, the {110} texture content is 26.4%, the {100} texture content is 9.7%, the yield strength is 989MPa, the tensile strength is 1295MPa, the elongation is 30.7%, the product of strength and ductility is 39.7Gpa·%, and the r value is 0.982.

[0330] According to the above results, it can be seen that the total content of {111} texture, the product of strength and ductility, and the r value of the steel material in the embodiment are significantly improved, and the cold stamping forming performance is obviously improved. The actual cold bending experimental results are as shown in Figure 45 Under the conditions of bending angle 0° and bending core diameter 16mm, the steel material processed by the process provided by the present application does not crack, which shows that the process provided by the present application can make the strength of the cold-formed steel material reach the strength of the general hot-formed steel, and the cold-forming performance is excellent.

[0331] In addition, the r values of two experimental steels (experimental steel A and experimental steel B) disclosed in the paper “Annealing Process Research of Rare Earth Micro-alloyed Low-cost Third Generation Automobile Steel” by Zhao Qingbo et al. in the background art are measured. Under the conditions of the best heat treatment process and deformation degree provided in the paper, the two experimental steels are annealed, and the measured r values are as shown in Table 14.

[0332] Table 14 r value measurement results of experimental steels

[0333]

[0334] As can be seen from the table, the r value of the test steel with 9ppm of Ce added after heat treatment is higher than that of the experimental steel without adding rare earth elements, and the r value is the highest under the condition of direct annealing in the two-phase region (645℃-15min), which is 0.958; and the r value of the experimental steel processed by the processing process of the embodiment, especially the experimental steel with a cold rolling deformation degree of 73% and direct annealing in the two-phase region at 640℃ for 15min, is 0.982, which is higher than the highest r value of the experimental steels in Table 14.

[0335] In the embodiment, the highest residual austenite content of the experimental steel without rare earth and the experimental steel with rare earth is 25.8% and 26.7% respectively, the highest tensile strength is 1230MPa and 1265MPa respectively, the highest elongation is 30.9% and 31.8% respectively, and the highest product of strength and elongation is 38GPa·% and 40.2GPa·% respectively. In the background art, the highest tensile strength of the medium manganese steel containing 9ppm of rare earth element is 1030MPa, the highest elongation is 31.94%, and the highest product of strength and elongation is 32.9GPa·%. The maximum product of strength and elongation of the experimental steel in the embodiment is 22.2% higher than that of the experimental steel containing 9ppm of rare earth element in the background art.

[0336] The above results show that the favorable texture content, product of strength and elongation, and r value of the experimental steel in the embodiment are improved compared with the experimental steel in the background art, and the experimental steel in the embodiment has better strength and plasticity, that is, better cold stamping forming performance.

Claims

1. A process for improving the cold press formability of 1000 MPa grade automotive steel, characterized in that, The method comprises the following steps: Step A: smelting and casting: the medium manganese steel raw material is placed in a vacuum induction furnace for smelting, and after the medium manganese steel raw material is completely melted, the rare earth iron alloy is added, and after the rare earth iron alloy is completely melted, homogenization treatment is performed, and then smelting and casting are performed to obtain a blank; the rare earth iron alloy is a Ce-Fe alloy containing 30wt% Ce; in step E, the mass fraction of Ce in the automobile steel is 0.0010wt%-0.0050wt%, the mass fraction of C is 0.11wt%-0.14wt%, the mass fraction of Mn is 5.0wt%-5.8wt%, the mass fraction of Si is 0.2wt%-0.3wt%, the mass fraction of Al is 0.020wt%-0.025wt%, the mass fraction of Cu is 0.24wt%-0.28wt%, the mass fraction of Ni is 0.24wt%-0.28wt%, the mass fraction of Nb is 0.014wt%-0.020wt%, the mass fraction of Ti is 0.018wt%-0.030wt%, the mass fraction of P is less than or equal to 0.020wt%, the mass fraction of S is less than or equal to 0.005wt%, the mass fraction of N is less than or equal to 0.01wt%, and the balance is Fe; Before smelting starts, the vacuum induction furnace is first pumped for 20-30min, during the smelting process, the vacuum degree in the vacuum induction furnace is less than or equal to 70Pa, the smelting temperature is 1600-1700℃, the holding time is 90-120min, and the homogenization treatment time is 5-10min; Step B: hot rolling: after the blank is heated and kept, hot rolling is performed to obtain a hot-rolled plate; The blank is heated to 1200-1250℃, and then kept for 4-5h for diffusion annealing; the opening rolling temperature range during hot rolling is 1170-1180℃, and the final rolling temperature is 860-880℃; the blank is sequentially rolled to a thickness of 110mm, 95mm, 67mm, 54mm, 48mm, 32mm, 18mm, 13mm, 8mm, 5mm and 3mm; Step C: pickling: after the hot-rolled plate is soaked in hydrochloric acid, the hydrochloric acid on the surface of the hot-rolled plate is washed away with clean water; then the hot-rolled plate is blown dry and oiled on the surface to obtain a pickled hot-rolled plate; Step D: cold rolling: the pickled hot-rolled plate is first annealed and softened, and then cold-rolled to obtain a cold-rolled plate; The annealing and softening temperature is 800-850℃, the annealing and softening time is 4-5h, and the cold rolling deformation degree is 33%-73%; Step E: annealing: the cold-rolled plate is annealed, and after the annealing treatment is completed, a 1000MPa grade automobile steel with excellent cold stamping forming performance is obtained; The annealing treatment is ART annealing or two-phase zone direct annealing; The ART annealing method is as follows: the cold-rolled plate is heated to 790-820 ℃ at a heating rate of 5-15 ℃ / min, and then water-quenched after being kept at 790-820 ℃ for 5 min, and the cooling rate is greater than 15 ℃ / min; then the cold-rolled plate is heated to 620-660 ℃ at a heating rate of 5-15 ℃ / min, and then air-cooled to room temperature after being kept at 620-660 ℃ for 5-25 min, and the cooling rate is 8-12 ℃ / min. The two-phase region direct annealing method is as follows: the cold-rolled plate is heated to 620-660 ℃ at a heating rate of 5-15 ℃ / min, and then air-cooled to room temperature after being kept at 620-660 ℃ for 5-25 min, and the cooling rate is 8-12 ℃ / min.

2. Process for improving the cold stamping formability of 1000 MPa grade automotive steel according to claim 1 characterized in that, In step C, the hot-rolled plate is immersed in concentrated hydrochloric acid with a pH of 0-1 for 10-12 h.

3. Process for improving the cold stamping formability of 1000 MPa grade automotive steel according to claim 1 characterized in that, The ART annealing method is as follows: the cold-rolled plate is heated to 800 ℃ at a heating rate of 10 ℃ / min, and then water-quenched after being kept at 800 ℃ for 5 min, and the cooling rate is 20 ℃ / min; then the cold-rolled plate is heated to 640 ℃ at a heating rate of 10 ℃ / min, and then air-cooled to room temperature after being kept at 640 ℃ for 15 min.

4. Process for improving the cold stamping formability of 1000 MPa grade automotive steel according to claim 1 characterized in that, The two-phase region direct annealing method is as follows: the cold-rolled plate is heated to 640 ℃ at a heating rate of 10 ℃ / min, and then air-cooled to room temperature after being kept at 640 ℃ for 15 min.

5. Process for improving the cold press forming properties of 1000 MPa grade automotive steel according to claim 1 characterized in that, In step A, the rare earth iron alloy is a Ce-Fe alloy containing 30 wt% of Ce; in step E, the mass fraction of Ce in the automobile steel is 0.0048 wt%, the mass fraction of C is 0.13 wt%, the mass fraction of Mn is 5.59 wt%, the mass fraction of Si is 0.3 wt%, the mass fraction of Al is 0.021 wt%, the mass fraction of Cu is 0.26 wt%, the mass fraction of Ni is 0.261 wt%, the mass fraction of Nb is 0.016 wt%, the mass fraction of Ti is 0.023 wt%, the mass fraction of P is 0.015 wt%, the mass fraction of S is 0.005 wt%, the mass fraction of N is 0.01 wt%, and the balance is Fe; In step A, before the smelting starts, the vacuum induction furnace is first evacuated for 20 min, and during the smelting process, the vacuum degree in the vacuum induction furnace is less than or equal to 70 Pa, the smelting temperature is 1600 ℃, the holding time is 90 min, and the homogenization treatment time is 5 min; In step B, the blank is heated to 1200 ℃, and then kept for 5 h for diffusion annealing; the opening rolling temperature is 1180 ℃, and the final rolling temperature is 860 ℃; the blank is sequentially rolled to a thickness of 110 mm, 95 mm, 67 mm, 54 mm, 48 mm, 32 mm, 18 mm, 13 mm, 8 mm, 5 mm and 3 mm; In step C, the hot-rolled plate is immersed in concentrated hydrochloric acid with a pH of 0 for 10 h; In step D, the annealing softening temperature is 850 ℃, the annealing softening time is 4 h, and the cold rolling deformation degree is 73%; In step E, the annealing treatment is ART annealing or two-phase region direct annealing; The method of ART annealing is: the cold-rolled plate is heated to 800℃ at a heating rate of 10℃ / min, and after being kept at 800℃ for 5min, water quenching is carried out, and the cooling rate is 20℃ / min; then the cold-rolled plate is heated to 640℃ at a heating rate of 10℃ / min, and after being kept at 640℃ for 15min, air cooling is carried out to room temperature, and the cooling rate of air cooling is 10℃ / min; The method of two-phase region direct annealing is: the cold-rolled plate is heated to 640℃ at a heating rate of 10℃ / min, and after being kept at 640℃ for 15min, air cooling is carried out to room temperature, and the cooling rate of air cooling is 10℃ / min.

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