Hot-rolled high-strength steel with excellent chambering and bending performance and manufacturing method of hot-rolled high-strength steel

Abstract

The invention discloses hot-rolled high-strength steel with excellent chambering and bending properties, which contains Fe and inevitable impurities, and further contains the following chemical elements in percentage by mass: 0.03-0.08% of C, 0.01-1.2% of Si, 1.2-1.8% of Mn, 0.0005-0.008% of S, 0.02-1.0% of Al, less than or equal to 0.0035% of B, less than or equal to 0.015% of P, less than or equal to 0.015% of S and the balance of Fe. And at least one of: 0.0002 to 0.01% of Mg, and 0.0002 to 0.01% of Ca; 0 < Ti < = 0.13%, 0 < Nb < = 0.06%, and 0 < V < = 0.20%; the grain boundary density of the microstructure of the hot-rolled high-strength steel is 0.6-2.0 [mu] m <-1 >. Correspondingly, the invention further provides a manufacturing method of the hot-rolled high-strength steel.

Description

Technical Field

[0001] The invention relates to a steel plate and a manufacturing method thereof, and in particular to a hot-rolled high-strength steel and a manufacturing method thereof. Background Art

[0002] With the development of automobile lightweight technology, the proportion of high-strength steel plates in automobile structural parts is increasing. At present, many models use 80kg-level steel plates to produce automobile chassis parts. Automobile chassis parts such as control arms have extremely high requirements for hole expansion performance in the forming process, including stamping, flanging, and hole expansion.

[0003] A Chinese patent document with publication number CN104513930A and publication date April 15, 2015, entitled "Ultra-high strength hot-rolled complex phase steel plate and strip with good bending and hole expansion performance and manufacturing method thereof", discloses a steel plate with a microstructure of at least 80% bainite, ferrite, martensite and residual austenite, a hole expansion rate of more than 50%, and a 180° cold bending of 0a.

[0004] The Chinese patent document with publication number CN105154769A and publication date December 16, 2015, entitled "A 780MPa grade hot-rolled high-strength high-hole expansion steel and its manufacturing method", discloses that the hot-rolled coiling temperature of the steel plate is 600-700°C, there is sufficient time for precipitation within the temperature range of 600-700°C, and the cooling rate after coiling is controlled to be ≤20°C / h, and its structure is ferrite.

[0005] The Chinese patent document with publication number CN112575267A and publication date March 30, 2021, and titled "A high hole expansion complex phase steel and its manufacturing method", discloses transverse tensile strength ≥780MPa, yield strength ≥700MPa, A50 elongation ≥15%, punching hole expansion rate ≥50%, and microstructure of bainite and ferrite.

[0006] It can be seen that the above patent documents do not involve the influence of grain boundary density on hole expansion rate and bending performance. Summary of the invention

[0007] One of the purposes of the present invention is to provide a hot-rolled high-strength steel with excellent hole expansion and bending performance and a manufacturing method thereof. The hot-rolled high-strength steel has high punching quality, excellent hole expansion rate and bending performance, and can be used as automobile body structural parts and automobile chassis parts.

[0008] In order to achieve the above object, the present invention provides a hot-rolled high-strength steel with excellent hole expansion and bending performance, which contains Fe and inevitable impurities, and further contains the following chemical elements with the following mass percentages:

[0009] C: 0.03-0.08%, Si: 0.01-1.2%, Mn: 1.2-1.8%, S: 0.0005-0.008%, Al: 0.02-1.0%, B ≤ 0.0035%; and

[0010] At least one of Mg: 0.0002-0.01% and Ca: 0.0002-0.01%;

[0011] At least one of 0<Ti≤0.13%, 0<Nb≤0.06%, 0<V≤0.20%;

[0012] The grain boundary density of the hot-rolled high-strength steel microstructure is 0.6-2.0 μm -1 .

[0013] Furthermore, in the hot-rolled high-strength steel of the present invention, the mass percentage of each chemical element is:

[0014] C: 0.03-0.08%, Si: 0.01-1.2%, Mn: 1.2-1.8%, S: 0.0005-0.008%, Al: 0.02-1.0%, B≤0.0035%; at least one of Mg: 0.0002-0.01% and Ca: 0.0002-0.01%; at least one of 0<Ti≤0.13%, 0<Nb≤0.06%, 0<V≤0.20%; the balance is Fe and other inevitable impurities.

[0015] In the hot-rolled high-strength steel of the present invention, the design principles of each chemical element are specifically described as follows:

[0016] C: In the hot-rolled high-strength steel described in the present invention, the carbon content largely determines the tensile strength level of the steel plate. Carbon can be used for solid solution strengthening, and to form sufficient precipitation strengthening phases with Ti elements, etc., to ensure the strength of the steel; but when the mass percentage of carbon is too high, the carbide particles will be coarse, and it will be easy to form too many carbides (pearlite, cementite), which is not conducive to hole expansion performance. In order to ensure that the steel grade can achieve high hole expansion and have good forming and welding properties under the strength, in the hot-rolled high-strength steel described in the present invention, the mass percentage of C is controlled to be between 0.03-0.08%.

[0017] Si: In the hot-rolled high-strength steel described in the present invention, silicon can play a role in solid solution strengthening, improve the strength of the steel plate, and at the same time, adding silicon can inhibit the formation of harmful carbides, increase the ferrite ratio, and help improve the elongation of the steel plate. However, too high silicon content in the steel can easily cause the surface defect of 2FeO-SiO2 iron oxide scale to form on the surface of the steel plate, which has an adverse effect on the surface quality. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of silicon is controlled to be between 0.01-1.2%.

[0018] Mn: In the hot-rolled high-strength steel described in the present invention, manganese is a solid solution strengthening element. Manganese can delay the pearlite transformation, improve the hardenability of steel and reduce the bainite transformation temperature, refine the organizational substructure of the steel, and ensure the acquisition of the lath substructure organization, which can ensure the tensile strength of the product while having good formability. When the mass percentage of manganese is too low, it will lead to insufficient strength. However, if the mass percentage of manganese is too high, the plasticity of the steel plate will be reduced, and centerline segregation will occur, which will damage the formability. And when the Mn content is too high, it is easy to combine with S to form MnS, which will promote separation when punching or cutting the steel plate strip, further damaging the formability. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of Mn can be controlled between 1.2-1.8%.

[0019] S: In the hot-rolled high-strength steel described in the present invention, the S element is a beneficial element, not an impurity. It can combine with Mg and Ca to form sulfides, can become the nucleus of (Ti, Nb)N and TiN, and is beneficial to the refinement of (Ti, Nb)N and TiN, thereby improving the punching and trimming quality and the hole expansion performance. However, when the S content is too high, S and Mn form MnS, which deteriorates the hole expansion performance. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of S can be controlled between 0.0005-0.008%.

[0020] Al: In the hot-rolled high-strength steel described in the present invention, Al is a deoxidizing element of steel, which can reduce oxide inclusions in steel, purify steel, and help improve the formability of steel plates. At the same time, aluminum, like silicon, will inhibit the formation of harmful carbides, increase the ferrite ratio, and help improve the elongation of steel plates. However, when the mass percentage of aluminum is too high, oxidation will occur, further affecting continuous casting production. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of Al can be controlled to 0.02-1.0%.

[0021] B: In the hot-rolled high-strength steel described in the present invention, boron is beneficial to expand the bainite phase region and ensure that the steel plate can obtain bainite structure during cooling after rolling, which significantly improves the strength and hardness of the steel. However, too much B element will cause too much large block martensite structure to appear in the steel plate, resulting in a decrease in the hole expansion rate and elongation of the steel. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of B can be controlled to B≤0.0035%.

[0022] Mg: In the hot-rolled high-strength steel described in the present invention, magnesium can form oxides and sulfides with oxygen and sulfur, respectively. Compared with steel without Mg, the generated Mg-based oxides and Mg-based sulfides have smaller precipitates of TiN and (Ti, Nb)N and are evenly dispersed, which is beneficial to the improvement of hole expansion performance. However, when Mg is less than 0.0002%, the effect is not sufficient. When it exceeds 0.01%, excessive oxides and sulfides will be generated, which will deteriorate the hole expansion performance. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of Mg can be controlled to be between 0.0002-0.01%.

[0023] Ca: In the hot-rolled high-strength steel described in the present invention, calcium and magnesium have similar effects and can improve the morphology of sulfides such as MnS, so that long strips of MnS and other sulfides become spherical MnS, which is beneficial to improve the morphology of inclusions, thereby reducing the adverse effects of long strips of sulfides on hole expansion forming performance. However, the addition of too much calcium will increase the amount of calcium oxide, which is not conducive to hole expansion performance. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of Ca can be controlled to be between 0.0002-0.01%.

[0024] Ti: In the hot-rolled high-strength steel described in the present invention, titanium is one of the important fine grain strengthening and precipitation strengthening elements. Ti can increase the recrystallization temperature and refine the grain size during the hot rolling process. At the same time, the combination of Ti and C has a very good strengthening effect. However, the mass percentage of Ti should not be too much, otherwise large-sized TiN will be easily formed, which is not conducive to the impact toughness of the steel. Based on this, in the hot-rolled high-strength steel described in the present invention, the upper limit of the mass percentage of Ti can be controlled to 0.13%. In some preferred embodiments, Ti can be controlled between 0.05-0.13%.

[0025] Nb: In the hot-rolled high-strength steel of the present invention, Nb is one of the important precipitation strengthening and fine grain strengthening elements, but when the mass percentage of Nb is higher than 0.06%, the strengthening effect of Nb is close to saturation, and the cost is relatively high. Therefore, in order to exert the beneficial effects of the Nb element and control the production cost at the same time, in the hot-rolled high-strength steel of the present invention, the upper limit of the mass percentage content of the Nb element can be controlled to 0.06%, and in some preferred embodiments, Nb can be controlled between 0.01-0.05%.

[0026] V: In the hot-rolled high-strength steel described in the present invention, V is one of the important precipitation strengthening and fine grain strengthening elements. It cooperates with niobium to play a good role in refining austenite grains and precipitation strengthening. It exists in the form of fine precipitation during cooling after rolling or after coiling, and precipitation strengthening is used to improve strength. Therefore, in the hot-rolled high-strength steel described in the present invention, the upper limit of the mass percentage of vanadium can be controlled to 0.20%, and in some preferred embodiments, V can be controlled between 0.05-0.20%.

[0027] Furthermore, in the hot-rolled high-strength steel of the present invention, each chemical element satisfies at least one of the following formulas:

[0028] Mg / 24+Ca / 40≥O / 16×0.83(Formula 1);

[0029] S / 32≤Mg / 24+Ca / 40-O / 16×0.83(Formula 2);

[0030] S / 32×Mn / 55≤8.0×10 -6 (Formula 3);

[0031] In the formula, each chemical element is substituted into the value before the percentage sign of its mass percentage.

[0032] In the present invention, the content of O that can be effectively combined with Mg and Ca is 83% of the total O content, so the content of added Mg and Ca preferably satisfies the formula Mg / 24+Ca / 40≥O / 16×0.83.

[0033] In the present invention, when the S content in the steel is high, MnS is easily formed, which is detrimental to the hole expansion performance and bending performance of the material. Therefore, in addition to fixing O, the added Mg and Ca need to further fix S to reduce the formation of MnS, so the Mg and Ca contents preferably satisfy the formula S / 32≤Mg / 24+Ca / 40-O / 16×0.83 at the same time.

[0034] In the present invention, since a large amount of Mn is added to the high-strength steel, when the S content is also high, MnS will precipitate at high temperature, inhibiting the formation of MgS and CaS, and reducing the hole expansion and bending properties. Therefore, the Mn and S contents preferably satisfy the formula S / 32×Mn / 55≤8.0×10 -6 .

[0035] Furthermore, in the hot-rolled high-strength steel of the present invention, when the corresponding elements are contained, the mass percentage thereof further satisfies:

[0036] Ti: 0.05-0.13%;

[0037] Nb: 0.01-0.05%;

[0038] V: 0.05-0.20%.

[0039] Furthermore, in the hot-rolled high-strength steel described in the present invention, the mass percentage of each chemical element satisfies: 0.5≤4×C / (3.3Nb+3.4V+Ti)≤2.25 (Formula 4), where each chemical element is substituted into the value before the percentage sign of its mass percentage.

[0040] In the slab heating process of the present invention, the growth of the original austenite grains is prevented by the precipitates of Nb, Ti, and V. In the hot rolling process of the present invention, (Nb, Ti, V)C is conducive to increasing the recrystallization temperature and further refining the austenite grains. The precipitated (Nb, Ti, V)(C, N) or (Nb, Ti, V)(Cr, Mo)(C, N) is conducive to refining the bainite of the phase transformation and a small amount of martensite grains. In the laminar cooling process of the present invention, the nano-scale precipitation of (Nb, Ti, V)(C, N) or (Nb, Ti, V)(Cr, Mo)(C, N) plays a strong precipitation strengthening effect; its precipitation on the bainite and ferrite matrix is ​​conducive to the improvement of strength, especially the strength improvement of the ferrite matrix, further reducing the strength difference between ferrite and bainite and martensite phases, which is conducive to the improvement of hole expansion rate, and refining the size of bainite, quasi-polygonal ferrite, polygonal ferrite, martensite and other grains in the microstructure, which is conducive to the improvement of hole expansion performance and bending performance. Therefore, in order to obtain a better precipitation strengthening effect, the content of Nb, V and Ti needs to meet the following requirements: 0.5≤4×C / (3.3Nb+3.4V+Ti)≤2.25.

[0041] Furthermore, the hot-rolled high-strength steel of the present invention further contains at least one of the following chemical elements: 0<Cr≤0.7%, 0<Mo≤0.25%.

[0042] In a preferred embodiment of the present invention, at least one of Cr and Mo may be added.

[0043] Cr: In the hot-rolled high-strength steel described in the present invention, the chromium element can inhibit the formation of pearlite, which is beneficial to the formation of bainite structure, and ultimately helps to improve the strength and hole expansion rate. When the chromium content is too low, the effect on the phase transformation curve is not significant, but when the mass percentage of Cr is too high, on the one hand, it leads to higher costs, and on the other hand, it is easy to produce more martensitic structure. Based on this, in the hot-rolled high-strength steel described in the present invention, the upper limit of the mass percentage of Cr can be controlled to 0.7%.

[0044] Mo: In the hot-rolled high-strength steel described in the present invention, the Mo element can not only inhibit the production of pearlite, but also facilitate the formation of bainite structure and a small amount of martensite austenite islands. In addition, the Mo element can also promote the transformation of bainite microstructure at a higher temperature, thereby allowing the steel to be coiled at a higher temperature. This higher coiling temperature provides sufficient precipitation kinetics, thereby stimulating significant precipitation strengthening. In the present invention, the Mo element also plays a role in the composite precipitation process with Nb and Ti, and it can also reduce the possibility of coarse particle size of the precipitated particles. However, it should be noted that the Mo content in the steel should not be too high. When the Mo content in the steel is too high, it will not only lead to an increase in the alloy cost, but also easily form more martensite and austenite, which is not conducive to the performance of the steel. Therefore, in order to give full play to the beneficial effects of the Mo element, in the hot-rolled high-strength steel described in the present invention, the upper limit of the mass percentage of the Mo element can be controlled to 0.25%.

[0045] Furthermore, in the hot-rolled high-strength steel of the present invention:

[0046] When Cr<0.40%, Mo content is 0.05%-0.25%;

[0047] When the Cr content is 0.40-0.7%, Mo is not added.

[0048] In the present invention, Cr and Mo play an important role in the formation of bainite. The purpose of adding an appropriate amount of Cr and / or Mo elements is to obtain bainite structure, smaller martensite and pearlite during hot rolling and coiling, so as to increase the grain boundary density of the microstructure, inhibit crack propagation during hole expansion, and further improve hole expansion performance and bending performance.

[0049] Furthermore, among the inevitable impurities of the hot-rolled high-strength steel of the present invention, P≤0.02%, N≤0.005%, and O≤0.0040%.

[0050] In the hot-rolled high-strength steel of the present invention, the inevitable impurities are mainly P, N and O. Under the condition that technical conditions permit, in order to obtain steel with better performance and better quality, the content of impurity elements in the steel should be reduced as much as possible. Among them:

[0051] P: When the P content is too high, it will cause deterioration of weldability, workability and toughness. Based on this, in some embodiments, the mass percentage of P can be controlled to P≤0.02%.

[0052] N: The lower its content, the better, but nitrogen is an inevitable element in the steelmaking process. Although its content is small, it is combined with strong carbide-forming elements such as Ti. Since TiN and (Ti, Nb)N are square, there is a large stress concentration between their sharp corners and the matrix. The stress concentration between TiN and (Ti, Nb)N and the matrix is ​​easy to form cracks, which also has a great impact on the fracture toughness and hole expansion performance. Based on this, in some embodiments of the present invention, the mass percentage of N can be controlled to N≤0.005%.

[0053] O: Oxygen is an inevitable element in the steelmaking process. For the present invention, the content of O in the steel will remain after deoxidation, and it is easy to form oxides. The inclusions themselves will not have a significant adverse effect on the performance of the steel plate. At the same time, Al2O3 is easy to become the nucleus of TiN and promote the growth of TiN. Based on this, in the hot-rolled high-strength steel described in the present invention, the mass percentage of O can be controlled to O≤0.0040%.

[0054] Furthermore, in the hot-rolled high-strength steel of the present invention, its microstructure includes bainite, quasi-polygonal ferrite, polygonal ferrite and martensite.

[0055] Furthermore, in the hot-rolled high-strength steel of the present invention, the area fraction of bainite+quasi-polygonal ferrite is 70-95%, the area fraction of polygonal ferrite is 3-30%, and the area fraction of martensite is ≤5%.

[0056] Furthermore, in the hot-rolled high-strength steel of the present invention, the grain size of bainite and quasi-polygonal ferrite is ≤4 μm, the grain size of polygonal ferrite is ≤5 μm, and the grain size of martensite is ≤3 μm.

[0057] In this embodiment, the grain sizes of bainite, quasi-polygonal ferrite, polygonal ferrite and martensite in the microstructure are refined, which is beneficial to further improve the hole expansion performance and bending performance.

[0058] Furthermore, in the hot-rolled high-strength steel of the present invention, the size of the precipitates is ≤8 μm.

[0059] Furthermore, in the hot-rolled high-strength steel of the present invention, the proportion of precipitates with a size of ≤3.0 μm in the total precipitates is ≥50%.

[0060] Furthermore, in the hot-rolled high-strength steel described in the present invention, the precipitates include MgO, CaO, MgS, CaS, TiN, and (Nb, Ti)N particles.

[0061] In the present invention, Mg and Ca are added to form Mg and Ca-based oxides and sulfides, which can play a nucleation role for TiN and (Ti, Nb)N precipitates. The Mg and Ca-based oxides and sulfides precipitated at low temperatures inhibit the growth of TiN and (Ti, Nb)N by competing with TiN and (Ti, Nb)N for precipitation, and also inhibit the growth of MnS, further reducing the generation of fine and uniform pores on the punched section, reducing stress concentration during hole expansion, and ultimately improving hole expansion performance and bending performance.

[0062] Furthermore, in the hot-rolled high-strength steel described in the present invention, its yield strength is ≥660 MPa, tensile strength is ≥780 MPa, elongation A50 is ≥15%, punching hole expansion ratio is ≥75%, and VDA bending angle is ≥120°.

[0063] In addition, another object of the present invention is to provide a method for manufacturing hot-rolled high-strength steel, and the hot-rolled high-strength steel plate manufactured by this method has high punching quality, excellent hole expansion rate and bending performance.

[0064] In order to achieve the above object, the present invention proposes a method for manufacturing hot-rolled high-strength steel with excellent hole expansion and bending performance, which comprises the steps of:

[0065] Smelting and casting: the superheat is controlled at 15-55℃;

[0066] Heating: The slab is heated to 1200-1300℃ and kept warm for 1-3h;

[0067] Rolling: The rough rolling outlet temperature is controlled at 1000-1080℃, the final rolling temperature is 840-950℃, and after rolling, the steel is cooled to a coiling temperature of 300-620℃ at an average cooling rate of ≥30℃ / s, and then coiled at a cooling rate of ≤15℃ / s to room temperature.

[0068] Pickling.

[0069] Furthermore, in the manufacturing method described in the present invention, a galvanizing step is also included after the pickling step, thereby obtaining products with different coatings such as electro-galvanizing, hot-dip galvanizing, zinc-aluminum-magnesium, etc.

[0070] In the manufacturing method described in the present invention, during the smelting and casting process, the central segregation and inclusion levels of the continuous casting billet can be controlled by controlling the superheat, secondary cooling water, and adopting appropriate soft pressure, and the inclusion grade can be controlled to be less than grade 1.5.

[0071] In the manufacturing method described in the present invention, the superheat is controlled to be 15-55°C in order to better control the size and precipitate density of MgO, CaO, MgS, CaS, (Nb, Ti)N, and TiN. Large, brittle, and sharp-edged (Nb, Ti)N and TiN will become potential crack sources and greatly reduce the hole expansion performance and bending performance of this type of steel. In addition, the superheat of steelmaking has a greater impact. The greater the superheat, the more conducive it is to the control of inclusions, but the greater the superheat, the more conducive it is to the growth of TiN. Therefore, in the steelmaking process, in conjunction with the addition of Ca and Mg, the present invention can control the superheat to be 15-55°C to obtain precipitates in the steel plate of MgO, CaO, MgS, CaS, TiN, (Nb, Ti)N, whose size is ≤8μm, and further the proportion of precipitates ≤3.0um is ≥50%, which can further improve the hole expansion performance and bending performance.

[0072] For micro-alloy steels containing Nb, Ti, V, etc., the heating temperature of the slab is particularly important for the performance. Nb, Ti, V will have a large amount of large-sized (Nb, Ti, V) (C, N) precipitation during the continuous casting process. During the heating process of the slab, it is necessary to dissolve the alloying elements such as Nb, V, and Ti as much as possible to ensure the nano-scale precipitation of micro-alloys such as Ti in the subsequent hot rolling and coiling process, so the heating temperature is set to ≥1200°C. When the temperature exceeds 1300°C, there will be a tendency for grain coarsening, which is not conducive to the toughness of the steel plate, and the iron oxide scale is thick, which is not conducive to the dephosphorization of the iron oxide scale. Therefore, in the manufacturing method described in the present invention, the heating temperature is set to 1200-1300°C.

[0073] In the present invention, the rough rolling temperature control in the hot rolling process has a greater impact on microalloys such as Nb, V, and Ti. Microalloys such as Nb, V, and Ti will have Ti carbides and carbonitrides precipitated during the relatively low rough rolling temperature and finishing rolling process. The precipitation size of this process is relatively large, which is not conducive to the improvement of the final strength. Therefore, the rough rolling outlet temperature is controlled to be 1000-1080°C. In addition, the present invention controls the final rolling temperature to be 840-950°C, which can be rolled in the non-recrystallized area for grain refinement.

[0074] In the present invention, in order to obtain a microstructure of bainite + quasi-polygonal ferrite + polygonal ferrite + martensite, especially a microstructure with a grain boundary density of 0.6-2.0 μm -1 The composition design and manufacturing process need to work together. From the perspective of composition, you can choose to add one or both of Cr and Mo. The addition of Cr and / or Mo has effectively suppressed the formation of pearlite, which is beneficial to the improvement of grain boundary density, and further improves the hole expansion performance and bending performance. However, when its content is excessive, it is also easier to form blocky secondary martensite. Therefore, the addition of Cr and / or Mo also needs to be coordinated with the hot rolling laminar cooling process.

[0075] The laminar cooling after rolling of the present invention has an important influence on the microstructure grain boundary density, as well as the volume fraction of bainite, quasi-polygonal ferrite, blocky ferrite and martensite phase transformation. In the present invention, in order to obtain a suitable microstructure ratio, it is necessary to strictly control a reasonable laminar cooling process and hot rolling coiling temperature. When the coiling temperature is too high, the secondary martensite and polygonal ferrite are large in size, which is not conducive to the improvement of hole expansion performance and bending performance; when the coiling temperature is too low, primary martensite structure may appear, and the hole expansion rate and elongation are low. Therefore, the coiling temperature can be controlled to between 300-620°C. Further, the coiling temperature can be controlled to 350-600°C.

[0076] After coiling, cooling to room temperature at a cooling rate of ≤15°C / s is beneficial to promote further transformation of bainite, reduce the proportion of martensite, and improve hole expansion and bending properties.

[0077] Furthermore, in the method for manufacturing hot-rolled high-strength steel of the present invention, in the hot rolling step, the rolling speed is controlled to be 7-13 m / s.

[0078] Furthermore, in the method for manufacturing hot-rolled high-strength steel of the present invention, in the hot rolling step, the total reduction ratio is controlled to be ≥80%, the total reduction ratio of finish rolling is ≥50%, and the reduction ratio of a single pass of final rolling is ≤15%.

[0079] Furthermore, in the manufacturing method of hot-rolled high-strength steel described in the present invention, laminar cooling adopts a segmented cooling process, wherein the first section is cooled to an intermediate point temperature of 610-750°C at an average cooling rate of ≥100°C / s, and air cooling is continued for 4-10s; then the second section is cooled to the coiling temperature at an average cooling rate of ≥30°C / s.

[0080] In the present invention, the laminar cooling process has an important influence on the microstructure ratio and the microalloy precipitation of Nb, V, and Ti at the same time. The temperature range where the precipitation temperature of Nb, V, and Ti is the strongest is between 610 and 750°C. The actual coiling temperature is lower than this temperature, in order to better exert the nano-scale precipitation strengthening effect of (Nb, Ti, V) (C, N) or (Nb, Ti, V) (Cr, Mo) (C, N). At the same time, controlling the temperature during air cooling to between 610 and 750°C is also conducive to the formation of quasi-polygonal ferrite. In the laminar cooling process, two-stage cooling is further adopted, wherein the midpoint temperature is 610-750°C, the air cooling time is 4-10s, and the average cooling rate of the second stage is ≥30°C / s. At the same time, the precipitation of (Nb, Ti, V) (C, N) or (Nb, Ti, V) (Cr, Mo) (C, N) is conducive to the improvement of the microstructure grain boundary density, and is also conducive to the improvement of hole expansion performance and bending performance.

[0081] In some preferred embodiments, the coiling temperature can be controlled between 350-600° C., which is beneficial to further control the transformation of bainite and quasi-polygonal ferrite and microalloy precipitation, and control the strength, hole expansion performance and bending performance of the steel plate.

[0082] The hot-rolled high-strength steel with excellent hole expansion and bending performance and the manufacturing method thereof according to the present invention have the following advantages and beneficial effects:

[0083] The hot-rolled high-strength steel with excellent hole expansion and bending performance described in the present invention adopts precise Mg and Ca addition and matching manufacturing process to control the microstructure grain boundary density, and also controls the type and proportion of the microstructure, the size and amount of the precipitates. Therefore, hot-rolled high-strength steel plates can be produced on traditional hot rolling production lines or low-carbon short-process production lines.

[0084] The hot-rolled high-strength steel with excellent hole expansion and bending performance described in the present invention has the characteristics of high-strength ultra-high hole expansion and high cold bending, and its yield strength is ≥660MPa, tensile strength is ≥780MPa, elongation A50 is ≥15%, punching hole expansion rate is ≥75%, VDA bending angle is ≥120°, and can be used as automobile body structural parts and automobile chassis parts, and can also be used in other application fields requiring high strength, weight reduction, etc. BRIEF DESCRIPTION OF THE DRAWINGS

[0085] Figure 1 A grain boundary density map of Example 1 of the present invention is shown.

[0086] Figure 2 The grain boundary density map of Comparative Example 3 is shown.

[0087] Figure 3 The effect of grain boundary density on hole expansion ratio is shown.

[0088] Figure 4 The effect of grain boundary density on bending performance is shown.

[0089] Figure 5 It shows the hole expansion cracking caused by MnS coarseness.

[0090] Figure 6 A larger number of (Nb, Ti)(C, N) precipitate particles in Comparative Example 10 are shown. DETAILED DESCRIPTION

[0091] The hot-rolled high-strength steel with excellent hole expansion and bending performance described in the present invention will be further explained and illustrated below in conjunction with the drawings in the specification and specific embodiments. However, such explanation and illustration do not constitute an improper limitation on the technical solution of the present invention.

[0092] Examples 1-14 and Comparative Examples 1-14

[0093] The hot-rolled high-strength steel with excellent hole expansion and bending performance in each embodiment of the present invention is prepared by the following steps:

[0094] (1) Smelting and casting to obtain slabs:

[0095] In some embodiments, smelting can be performed by currently commonly used means. For example, a converter is used for steelmaking, and the molten steel is subjected to RH vacuum degassing treatment and LF furnace desulfurization treatment. The superheat degree of the steelmaking process is controlled to be 15-55°C. Table 1-1, Table 1-2 and Table 1-3 list the chemical element composition content and synergistic relationship of each embodiment of the present invention and the comparative example.

[0096] (2) Heat the slab to 1200-1300℃ and keep it at this temperature for 1-3h.

[0097] (3) Hot rolling: After rolling, the steel is cooled to a coiling temperature of 300-620°C at an average cooling rate of ≥30°C / s and then coiled. After coiling, the steel is cooled to room temperature at a rate of ≤15°C / s.

[0098] In some embodiments, laminar cooling adopts a segmented cooling process, wherein the first segment is cooled to an intermediate point temperature of 610-750°C at an average cooling rate of ≥100°C / s, and air cooling is continued for 4-10s; and then the second segment is cooled to the coiling temperature at an average cooling rate of ≥30°C / s.

[0099] In some embodiments, the rough rolling outlet temperature is controlled to be 1000-1080° C.; the final rolling temperature is 840-950° C., and the rolling speed is 7-13 m / s.

[0100] In some embodiments, the total reduction ratio is controlled to be ≥80%, the total reduction ratio of the finishing rolling is ≥50%, and the reduction ratio of a single pass of the final rolling is ≤15%.

[0101] (4) Pickling:

[0102] In some embodiments, the pickling tensile elongation can be controlled to be 0.2-1.8%; the pickling speed is 60-150 m / min, the temperature of the last pickling tank in the pickling process can be 80-90° C., and the iron ion concentration can be 30-40 g / L.

[0103] For the hot-rolled high-strength steel sheet obtained by the present invention, a galvanizing step can be further performed after pickling, thereby obtaining products with different coatings such as electro-galvanizing, hot-dip galvanizing, zinc-aluminum-magnesium, etc.

[0104] Table 1-1. (wt%, the balance is Fe and other inevitable impurities except P, O and N)

[0105]

[0106]

[0107] Table 1-2. (wt%, the balance is Fe and other inevitable impurities except P, O and N)

[0108] Steel Type Ti Nb V Cr Mo A 0.13 0 0 0.7 0 B 0 0 0.2 0.38 0.05 C 0.07 0.05 0 0.4 0 D 0.05 0 0.05 0 0 E 0 0.06 0 0 0 F 0.1 0 0 0 0.25 G 0.1 0 0.06 0 0.13 H 0 0.01 0.1 0.55 0 I 0.07 0.03 0 0.4 0 J 0.07 0.03 0 0.4 0 K 0.07 0.03 0 0.4 0 L 0.14 0 0 0.4 0 M 0.07 0.03 0 0.4 0 N 0.07 0.03 0 0.4 0

[0109] Table 1-3. (wt%, the balance is Fe and other inevitable impurities except P, O and N)

[0110]

[0111] Note: Steel types AH in Table 1-1, Table 1-2 and Table 1-3 are steel types used in the embodiments of the present invention, and steel type IN is steel type used in the comparative example. Each chemical element in each formula in Table 1-3 is substituted with the value before the percentage sign of its mass percentage content.

[0112] Table 2-1 and Table 2-2 list the specific process parameters of each embodiment of the present invention and the comparative example in the above process steps.

[0113] Table 2-1.

[0114]

[0115]

[0116]

[0117] Table 2-2.

[0118]

[0119]

[0120] The hot-rolled high-strength steels with excellent hole expansion and bending performance of Examples 1-14 and Comparative Examples 1-14 were sampled respectively, and the grain boundary density of the materials was analyzed by EBSD and TEM methods. The parameters used in the EBSD experiment were: acceleration voltage 20 kV, electron beam intensity 18 μA, working distance 16 mm, step length 0.1 μm, and tilt angle 70 degrees. After the EBSD experiment was completed, the obtained crystallographic data was processed using HKLChannel 5 software to analyze the grain boundary density distribution of the samples, and the test results are listed in Table 3.

[0121] In addition, the sample is electrolyzed to remove the precipitated phase in the sample or the precipitated phase is removed by extraction replication, and then the precipitated phase is photographed and single-point energy spectrum analyzed by transmission electron microscopy. Specifically, the physical and chemical phase analysis method uses a mixed aqueous solution of 5% potassium chloride + 1% citric acid, and the current density is controlled at 0.05A / cm2. The precipitated phase morphology, size, composition, etc. of samples with different processes are characterized and analyzed by transmission electron microscopy to determine the type and quantity of precipitates, and the test results are listed in Table 3.

[0122] Table 3.

[0123]

[0124]

[0125]

[0126] In addition, the mechanical properties of each embodiment and comparative sample were tested according to GB / T228.1-2010 standard, and the yield and tensile strengths were measured by taking JIS 5# tensile specimens along the longitudinal direction. The hole expansion rate test method was carried out according to the hole expansion rate test method specified in GBT24524-2021 standard. The specimen with a central hole was pressed into the concave die with a punch to expand the central hole of the specimen until necking or through cracks appeared at the edge of the plate hole. The bending performance was carried out according to VDA238-100 standard, and the bending performance was determined by the size of the bending angle.

[0127] Table 4.

[0128]

[0129]

[0130] As shown in Table 3 and Table 4, Examples 1-14 of the present invention obtain ideal organizational characteristics through reasonable chemical element composition design and optimized process parameters, wherein the precipitates of the example steel plate are MgO, CaO, MgS, CaS, TiN, (Nb, Ti) N, so that the steel plate obtains good performance. The yield strength finally obtained in each example of the present invention is greater than 660MPa, the tensile strength is greater than 780MPa, the elongation A50 is ≥15%, the punching hole expansion rate is ≥75%, and the VDA bending angle is greater than 120°. The particles are ≤8μm, and the proportion of precipitates ≤3.0um is ≥50%.

[0131] In addition, as shown in Table 3, the grain boundary density of each embodiment of the present invention is between 0.6 and 2.0 μm. -1 In addition, Figure 1 A grain boundary density map of Example 1 of the present invention is also shown. Figure 2 The grain boundary density map of Comparative Example 3 is shown.

[0132] The grain boundary density of the hot-rolled high-strength steel plate of the present invention has a direct influence on the punching and hole expansion performance and the bending performance.

[0133] Figure 3 The effect of grain boundary density on hole expansion ratio is shown. Figure 3 As shown, the grain boundary density is 0.6-2.0μm -1 Within the range, the greater the grain boundary density, the higher the hole expansion rate.

[0134] Figure 4 The effect of grain boundary density on bending performance is shown. Figure 4 As shown, the grain boundary density is 0.6-2.0μm -1 Within the range, the greater the grain boundary density, the greater the bending angle.

[0135] Unlike the present invention, the comparative examples cannot achieve the technical effects that the present invention can achieve because the ingredients or process parameters do not meet the design requirements of the present invention, wherein:

[0136] In Comparative Example 1, since the cooling speed after hot rolling and coiling is too fast, the proportion of martensitic phase transformation of supercooled austenite after coiling is too large, and the proportion of martensitic grains is too large, resulting in low hole expansion rate and bending performance.

[0137] Comparative Example 2 uses a lower coiling temperature, resulting in an excessively high primary martensite content in its microstructure, an excessively large martensite grain size, and an excessively large grain boundary density in the microstructure, which ultimately leads to a low hole expansion rate and bending performance, as well as a low elongation.

[0138] Comparative Example 3 uses a higher coiling temperature, resulting in a higher ferrite content and coarser grains in its microstructure. The area fraction of polygonal ferrite reaches 48%, and the grain size reaches 6.8 μm. At the same time, the grain size of bainite is 5.5 μm, which is relatively large. Figure 2 As shown, the grain boundary density of its microstructure is only 0.5um -1 , resulting in insufficient final strength, insufficient hole expansion rate and bending performance.

[0139] In Comparative Example 4, due to the high temperature at the middle point, the supercooling degree in the air cooling stage during the laminar cooling process is insufficient, and there is no ferrite phase transformation, which ultimately leads to a relatively low elongation of 14.5%.

[0140] In Comparative Example 5, due to the low rough rolling outlet temperature and the low finishing rolling temperature, polygonal ferrite phase transformation with larger size and higher area fraction occurs during the rolling process, and the grain boundary density is too small, which ultimately leads to lower yield strength and tensile strength.

[0141] In Comparative Example 6, due to the low heating temperature, the V content is not fully dissolved, and the coarse V (C, N) particles are not completely dissolved during the continuous casting process, which contributes little to the strength and also shows that the tensile strength of the steel plate is insufficient.

[0142] Since the superheat of Comparative Examples 7 and 8 cannot be controlled within a specific range, the superheat is low, which is not conducive to the floating of precipitated particles such as MgS, CaS, TiN, and is not conducive to the removal of inclusions. The superheat is too high, which is conducive to the growth of precipitated particles such as MgS, CaS, TiN, etc., which ultimately leads to larger sizes of precipitated particles such as MgS, CaS, TiN in the steel, and a larger proportion of particles with a size of ≤3.0μm. The particle sizes of Comparative Examples 7 and 8 are 12μm and 9μm, respectively, and the proportion of particles with a size of ≤3.0u is only 40% and 45%, respectively. Therefore, they also show lower hole expansion rate and bending performance.

[0143] In comparative example 9, due to the high S content, the MnS size in the final organization reaches 12 μm, which is relatively large, and the proportion of size ≤ 3.0 μm is only 40%, which has a great influence on the final hole expansion rate and bending performance, and the hole expansion rate and bending performance are insufficient. Figure 5 It shows the hole expansion cracking caused by the coarse MnS, and during the hole expansion process, MnS also exhibits brittle fracture.

[0144] In Comparative Example 10, due to the low content of Mg and Ca, the purpose of refining the precipitates cannot be achieved, resulting in the size of (Nb, Ti) (C, N) or (Nb, Ti, V) (Cr) (C, N) in the final precipitate particles reaching 9 μm, which is relatively large. The proportion of size ≤ 3.0 μm is only 40%, which has a great impact on the final hole expansion rate and bending performance, and the hole expansion rate and bending performance only reach 48% and 116°. Figure 6 A larger number of (Nb, Ti)(C, N) precipitate particles in Comparative Example 10 are shown.

[0145] In Comparative Example 11, due to the excessively high O content, there is an excess of CaO in the final precipitate particles, and the size of CaO is relatively large, which cannot serve the purpose of refining the precipitates. At the same time, the O content is relatively high, and the formed Al2O3 easily becomes the nucleus of TiN, promoting the growth of TiN, resulting in the size of (Nb, Ti)(C, N) or (Nb, Ti, V)(Cr)(C, N) in the final precipitate particles reaching 11μm, which is relatively large. The proportion of size ≤3.0μm is only 35%, which has a great influence on the final hole expansion rate and bending performance, and the hole expansion rate and bending performance are insufficient.

[0146] In Comparative Example 12, due to the high addition of C and Ti, on the one hand, the size of (Nb, Ti) (C, N) or (Nb, Ti, V) (Cr) (C, N) is larger and the number is larger; on the other hand, it leads to relatively high strength, which has a greater impact on the final hole expansion rate and bending performance, and the hole expansion rate and bending performance are insufficient.

[0147] In Comparative Example 13, the yield strength and tensile strength are insufficient due to the low addition of Mn.

[0148] In Comparative Example 14, the yield strength and tensile strength are insufficient due to the low addition of C.

[0149] It should be noted that the combination of the various technical features in this case is not limited to the combination described in the claims of this case or the combination described in the specific embodiments. All technical features recorded in this case can be freely combined or combined in any way unless there is a contradiction between them.

[0150] It should also be noted that the above-listed embodiments are only specific embodiments of the present invention. Obviously, the present invention is not limited to the above embodiments, and similar changes or modifications made therewith can be directly derived or easily associated with the contents disclosed by those skilled in the art from the present invention, and all should belong to the protection scope of the present invention.

Claims

1. A hot-rolled high-strength steel having excellent hole expansion and bending properties, containing Fe and inevitable impurities, characterized in that: It also contains the following chemical elements in the following mass percentages: C: 0.03-0.08%, Si: 0.01-1.2%, Mn: 1.2-1.8%, S: 0.0005-0.008%, Al: 0.02-1.0%, B≤0.0035%; At least one of Mg: 0.0002-0.01% and Ca: 0.0002-0.01%; At least one of 0<Ti≤0.13%, 0<Nb≤0.06%, 0<V≤0.20%; The grain boundary density of the hot-rolled high-strength steel microstructure is 0.6-2.0 μm -1 .

2. The hot-rolled high-strength steel according to claim 1, characterized in that: The mass percentage of each chemical element is: C: 0.03-0.08%, Si: 0.01-1.2%, Mn: 1.2-1.8%, S: 0.0005-0.008%, Al: 0.02-1.0%, B≤0.0035%; At least one of Mg: 0.0002-0.01% and Ca: 0.0002-0.01%; At least one of 0<Ti≤0.13%, 0<Nb≤0.06%, 0<V≤0.20%; the balance is Fe and other inevitable impurities.

3. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: Each chemical element satisfies at least one of the following formulas: Mg / 24+Ca / 40≥O / 16×0.83; S / 32≤Mg / 24+Ca / 40-O / 16×0.83; S / 32×Mn / 55≤8.0×10 -6 ; In the formula, each chemical element is substituted into the value before the percentage sign of its mass percentage.

4. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: When the corresponding elements are contained, their mass percentage further satisfies: Ti: 0.05-0.13%; Nb: 0.01-0.05%; V：0.05-0.20％。 5. The hot-rolled high-strength steel according to claim 4, characterized in that: The mass percentage of each chemical element satisfies: 0.5≤4×C / (3.3Nb+3.4V+Ti)≤2.25, where each chemical element is substituted into the value before the percentage sign of its mass percentage.

6. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: It also contains at least one of the following chemical elements: 0<Cr≤0.7%, 0<Mo≤0.25%.

7. The hot-rolled high-strength steel according to claim 6, characterized in that: When Cr<0.40%, Mo content is 0.05%-0.25%; When the Cr content is 0.40-0.7%, Mo is not added.

8. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: Among the inevitable impurities, P ≤ 0.02%, N ≤ 0.005%, O ≤ 0.0040%.

9. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: Its microstructure includes bainite, quasi-polygonal ferrite, polygonal ferrite and martensite.

10. The hot-rolled high-strength steel according to claim 9, characterized in that: The area fraction of bainite+quasi-polygonal ferrite is 70-95%, the area fraction of polygonal ferrite is 3-30%, and the area fraction of martensite is ≤5%.

11. The hot-rolled high-strength steel according to claim 9, characterized in that: The grain size of bainite and quasi-polygonal ferrite is ≤4μm, the grain size of polygonal ferrite is ≤5μm, and the grain size of martensite is ≤3μm.

12. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: The size of the precipitates is ≤8μm.

13. The hot-rolled high-strength steel according to claim 12, characterized in that: The proportion of precipitates with a size of ≤3.0 μm in the total precipitates is ≥50%.

14. The hot-rolled high-strength steel according to claim 12, characterized in that: The precipitate includes MgO, CaO, MgS, CaS, TiN, (Nb, Ti)N particles.

15. The hot-rolled high-strength steel according to claim 1 or 2, characterized in that: Its yield strength ≥660MPa, tensile strength ≥780MPa, elongation A50 ≥15%, punching expansion rate ≥75%, VDA bending angle ≥120°.

16. The method for producing hot-rolled high-strength steel according to any one of claims 1 to 15, characterized in that: It includes the steps of: Smelting and casting: the superheat is controlled at 15-55℃; Heating: The slab is heated to 1200-1300℃ and kept warm for 1-3h; Rolling: The rough rolling outlet temperature is controlled at 1000-1080℃, the final rolling temperature is 840-950℃, and after rolling, the steel is cooled to a coiling temperature of 300-620℃ at an average cooling rate of ≥30℃ / s, and then coiled at a cooling rate of ≤15℃ / s to room temperature. Pickling.

17. The manufacturing method according to claim 16, characterized in that: In the hot rolling step, the rolling speed is 7-13 m / s.

18. The manufacturing method according to claim 16, characterized in that: In the hot rolling step, the total reduction ratio is controlled to be ≥80%, the total reduction ratio of the finishing rolling is ≥50%, and the reduction ratio of a single pass of the final rolling is ≤15%.

19. The manufacturing method according to claim 16, characterized in that: The coiling temperature is 350-600°C.

20. The manufacturing method according to claim 16, characterized in that: Galvanizing is also included after the pickling step.

21. The manufacturing method according to any one of claims 16 to 20, characterized in that: Laminar cooling adopts a segmented cooling process, in which the first section is cooled to the midpoint temperature of 610-750°C at an average cooling rate of ≥100°C / s, and then air-cooled for 4-10s; then the second section is cooled to the coiling temperature at an average cooling rate of ≥30°C / s.

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