WO2017219940A1 - 一种980MPa级热轧双相钢及其制造方法 - Google Patents

一种980MPa级热轧双相钢及其制造方法 Download PDF

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WO2017219940A1
WO2017219940A1 PCT/CN2017/088965 CN2017088965W WO2017219940A1 WO 2017219940 A1 WO2017219940 A1 WO 2017219940A1 CN 2017088965 W CN2017088965 W CN 2017088965W WO 2017219940 A1 WO2017219940 A1 WO 2017219940A1
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steel
hot
rolled
mpa
ferrite
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French (fr)
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王焕荣
杨阿娜
张建苏
王巍
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Baoshan Iron and Steel Co Ltd
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Baoshan Iron and Steel Co Ltd
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Priority to JP2018566841A priority Critical patent/JP6797942B2/ja
Priority to US16/311,376 priority patent/US11268163B2/en
Priority to KR1020187038132A priority patent/KR102236345B1/ko
Publication of WO2017219940A1 publication Critical patent/WO2017219940A1/zh
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/74—Temperature control, e.g. by cooling or heating the rolls or the product
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention belongs to the field of hot rolled high strength steel, and particularly relates to a 980 MPa grade hot rolled dual phase steel and a manufacturing method thereof.
  • the steel for wheel of commercial vehicles, especially heavy trucks, is usually manufactured using duplex steel.
  • the wheels of some economical cars also use steel wheels to reduce costs.
  • the use of high-strength dual-phase steel to manufacture wheels can effectively reduce the weight of the wheel.
  • DP600 ie tensile strength 600MPa dual-phase steel
  • the 780MPa grade DP780 dual phase steel can further reduce the wheel weight by about 5-10%.
  • most of the wheel factories used in China have low-intensity dual-phase steels of less than 600 MPa, and higher-strength dual-phase steels such as DP780 are not used much.
  • duplex steel The reason why automobile wheels are widely used in duplex steel is mainly due to the low yield strength, high tensile strength, low yield ratio, continuous yielding and good formability of duplex steel itself. Porosity and so on. From the trend of industry development, the overall strength of steel for wheels is still developing toward higher strength, and the strength of wheel steels currently used is generally not high, between 500 and 600 MPa. With the increasingly strict national environmental laws and regulations and the implementation of national emission restrictions, in addition to passenger cars, the lightweight of cars in the commercial vehicle sector has also become a trend of industry development. In response to this development trend, it is necessary to develop higher-grade hot-rolled dual-phase steels such as 780 MPa and 980 MPa to meet the future development needs of wheel steel. In addition, high-strength dual-phase steel can also be applied to some automotive structural parts, such as automotive girders, anti-collision beams and so on.
  • high-strength dual-phase steels are mainly divided into two categories: one is cold-rolled high-strength dual-phase steel; the other is hot-rolled high-strength dual-phase steel.
  • the strength of cold-rolled high-strength dual-phase steel has reached 1180 MPa, while the strength of hot-rolled high-strength dual-phase steel has not reached 780 MPa or more.
  • the difference between cold-rolled and hot-rolled dual-phase steels is that the demand for high-strength dual-phase steel is mainly for cold-rolled products, and the demand for hot-rolled high-strength dual-phase steel is not strong; secondly, hot rolling production line and There are big differences in the cold rolling line, many of them are cold It is difficult to achieve the products that can be produced on the rolling line on the hot rolling line, and it is necessary to redesign and optimize the composition and process of the steel. As mentioned above, with the continuous development of the industry, the demand for hot-rolled high-strength dual-phase steel is bound to increase.
  • the object of the present invention is to provide a 980 MPa grade hot-rolled dual-phase steel and a manufacturing method thereof.
  • the hot-rolled dual-phase steel has a yield strength ⁇ 500 MPa, a tensile strength ⁇ 980 MPa, an elongation A 80 ⁇ 12%, and exhibits excellent performance.
  • Strength, plasticity and toughness can be applied to parts such as wheels that require good formability and high strength and thinning.
  • the invention adds a higher content of Si to ensure a certain amount of ferrite structure in a limited hot rolling air cooling time and expands the process window for ferrite formation; the main purpose of the composite addition of Nb and Ti is the largest in the finishing rolling stage.
  • the austenite grains are refined to a greater extent, so that the ferrite formed after the phase transformation is finer, which is advantageous for improving the strength and plasticity of the steel sheet.
  • the invention can obtain high-strength ferritic martensitic dual-phase steel with yield strength ⁇ 500MPa and tensile strength ⁇ 980MPa by precise control of ferrite and martensite content in the structure.
  • a 980MPa grade hot-rolled dual-phase steel having a chemical composition weight percentage of C: 0.10 to 0.20%, Si: 0.8 to 2.0%, Mn: 1.0 to 2.0%, P ⁇ 0.02%, S ⁇ 0.005%, O ⁇ 0.003%, Al: 0.02 to 0.06%, N ⁇ 0.006%, Nb: 0.01 to 0.06%, Ti: 0.01 to 0.05%, the balance being Fe and unavoidable impurities, and the above elements simultaneously satisfy the following relationship: 0.05% ⁇ Nb + Ti ⁇ 0.10%.
  • the chemical composition of the hot rolled duplex steel Si: 1.2 to 1.8% by weight.
  • Nb 0.03 to 0.05% by weight.
  • the chemical composition of the hot rolled duplex steel Ti: 0.02 to 0.04%, and the weight is 100 The ratio is calculated.
  • the microstructure of the hot-rolled dual-phase steel is fine ferrite + martensite, the volume fraction of ferrite is 20 to 35%, and the average grain size of ferrite is 5 to 10 ⁇ m; The volume fraction of the body is 65-80%, and the equivalent grain size of the martensitic ferrite is 15-20 ⁇ m.
  • the hot-rolled dual-phase steel of the invention has a yield strength ⁇ 500 MPa, a tensile strength ⁇ 980 MPa, and an elongation A 80 ⁇ 12%.
  • Carbon is a basic element in steel and is one of the important elements in the present invention. Carbon expands the austenite phase region and stabilizes austenite. Carbon as a gap atom in steel plays a very important role in improving the strength of steel, and has the greatest influence on the yield strength and tensile strength of steel.
  • the carbon content of the steel of the present invention must be controlled between 0.1 and 0.2%, preferably in the range of 0.14 to 0.18%.
  • Silicon is a basic element in steel and is also an important element in the present invention. This is because to obtain high-strength dual-phase steel with a tensile strength of 980 MPa or more, on the one hand, it is necessary to control the size and quantity of ferrite, and at the same time, to increase the strength of martensite, which requires proper improvement of carbon and composition in the composition design.
  • the content of manganese Both carbon and manganese are elements that expand the austenite region and stabilize austenite. It is difficult to form a sufficient amount of ferrite in a short time (usually ⁇ 10s) during hot rolling air cooling, which requires higher addition.
  • the content of silicon is a basic element in steel and is also an important element in the present invention. This is because to obtain high-strength dual-phase steel with a tensile strength of 980 MPa or more, on the one hand, it is necessary to control the size and quantity of ferrite, and at the same time, to increase the strength of martensite, which requires proper
  • the addition of silicon can significantly promote the formation of ferrite, enlarge the process window of ferrite formation, purify ferrite, and also play a partial strengthening role.
  • This effect of silicon must be exhibited when the content thereof is 0.8% or more, but the content of Si is not too high, otherwise the impact toughness of the steel sheet after rolling is deteriorated. Therefore, the silicon content in the steel of the present invention is controlled to be between 0.8 and 2.0%, preferably in the range of 1.2 to 1.8%.
  • Manganese is also the most basic element in steel and is one of the most important elements in the present invention. It is well known that manganese is an important element in expanding the austenite phase region, which can reduce the critical quenching speed of steel, stabilize austenite, refine grains, and delay the transformation of austenite to pearlite.
  • the manganese content in order to ensure the strength of the steel sheet, the manganese content should be controlled above 1.0%, the manganese content is too low, the supercooled austenite is not stable enough, and it is easy to be converted into a pearlite type structure when air-cooled; at the same time, the manganese content is also It should not exceed 2.0%.
  • the Mn content in the steel of the present invention is controlled to be 1.0 to 2.0%, preferably in the range of 1.4 to 1.8%.
  • Phosphorus is an impurity element in steel. It is easy to be segregated to the grain boundary. When the content of phosphorus in the steel is high ( ⁇ 0.1%), Fe 2 P is formed to precipitate around the grain, which reduces the plasticity and toughness of the steel. Therefore, the lower the content, the better. It is better within 0.02% and does not increase the cost of steel making.
  • Sulfur is an impurity element in steel. Sulfur in steel usually combines with manganese to form MnS inclusions. Especially when the content of sulfur and manganese is high, more MnS will be formed in the steel, and MnS itself has certain plasticity. MnS along the subsequent rolling process The rolling direction is deformed to reduce the transverse tensile properties of the steel sheet. Therefore, the lower the sulfur content in steel, the better, the actual production is usually controlled within 0.005%.
  • Aluminum is another important alloying element in steel except C, Si, Mn, P and S.
  • the basic role of aluminum in the present invention is primarily to deoxidize during the steelmaking process.
  • the content of aluminum in the steel is generally not less than 0.02%; meanwhile, if the content of aluminum exceeds 0.06%, the effect of refining the grains is weakened.
  • the present invention can control the content of aluminum in the steel to 0.02-0.06%.
  • Nitrogen is an impurity element in the present invention, and the lower the content, the better. Nitrogen is also an inevitable element in steel. Usually, if no special control is carried out during the steel making process, the residual content of nitrogen in the steel is usually ⁇ 0.006%. These solid solution or free nitrogen elements must be fixed by forming a certain nitride. Otherwise, the free nitrogen atom is very unfavorable to the impact toughness of steel, and it is easy to form a full-length sawtooth crack defect during strip rolling. . In the present invention, by adding a titanium element, it is combined with nitrogen to form a stable TiN to fix a nitrogen atom. Therefore, the nitrogen content in the steel of the present invention is controlled within 0.006% and the lower the better.
  • ⁇ is also one of the key elements in the present invention. Since hot-rolled dual-phase steel of 980 MPa and above usually needs to add higher silicon to promote the formation of ferrite phase in the rolling air cooling section, the addition of high silicon generally improves the brittleness of martensite. In the present invention, although the content of carbon itself is ⁇ 0.20%, after a certain amount of ferrite is precipitated, the carbon atoms in the ferrite are discharged into the untransformed austenite, so that the carbon content in the remaining austenite is rich. As a result, the final formed martensite has a higher actual carbon content and a larger martensite brittleness, which is further exacerbated by the addition of high silicon.
  • the low-temperature impact toughness of high-Si hot-rolled dual-phase steel is usually low.
  • a small amount of niobium is added to the alloy composition design, which can be effectively improved by refining the crystal grains. Impact toughness of duplex steel.
  • the addition of bismuth plays two roles: first, in the high temperature stage, solid solution enthalpy plays a role in solute dragging on austenite grain growth; second, in the finishing rolling stage, pinning austenite through bismuth carbonitride
  • the grain boundary refines the austenite grains, and refines the final transformed ferrite and martensite to improve the impact toughness of the duplex steel. Therefore, the niobium content in the steel of the present invention is controlled to be 0.01 to 0.06%, preferably in the range of 0.03 to 0.05%.
  • Titanium is one of the important elements in the present invention. Titanium plays two main roles in the present invention. First, it combines with the impurity element nitrogen in steel to form TiN, which plays a nitrogen-fixing effect. Second, it cooperates with niobium to optimize the role of refining austenite grains.
  • the free nitrogen atom in steel is very unfavorable to the impact toughness of steel.
  • the addition of trace titanium can fix the free nitrogen.
  • the content of titanium should not be too much, otherwise it is easy to form TiN with large size and impact toughness of steel. It is also unfavorable; the test proves that only Nb is added to the steel without adding Ti.
  • the continuous casting billet is prone to corner cracking, and the addition of trace titanium can effectively improve the corner cracking problem; meanwhile, in the present invention, Nb and Ti As long as the content is controlled within the range of 0.05% ⁇ Nb + Ti ⁇ 0.10%, the fine grain effect can be achieved and the cost is low. Therefore, the titanium content in the steel of the present invention is controlled in the range of 0.01 to 0.05%, preferably in the range of 0.02 to 0.04%.
  • Oxygen is an inevitable element in the steel making process.
  • the oxygen content in the steel can generally reach 30 ppm or less after deoxidation by aluminum, and does not cause significant adverse effects on the performance of the steel sheet. Therefore, the oxygen content in the steel can be controlled within 30 ppm.
  • the method for manufacturing 980 MPa grade hot-rolled dual-phase steel according to the present invention comprises the following steps:
  • the rolling temperature is 1030 ⁇ 1150°C, 3 ⁇ 5 passes rough rolling is performed at 1000°C or above and the cumulative deformation is ⁇ 50%; the intermediate billet temperature is 900 ⁇ 950°C, and then 3 ⁇ 5 passes are refined.
  • Rolling and cumulative deformation ⁇ 70%; finishing rolling temperature is 800-900 °C, after the end rolling, the steel plate is cooled to 600-700 °C at a cooling rate of ⁇ 100 °C/s; after air cooling for 5-10 seconds, the steel plate is further Quench at a cooling rate of 30 to 50 ° C / s to ⁇ 200 ° C, and after coiling, cool to room temperature at a cooling rate of ⁇ 20 ° C / h.
  • the rhythm of the rolling process should be completed as quickly as possible during the rough rolling and finishing rolling stages. After the end of the final rolling, it should be rapidly cooled to the intermediate cooling temperature at a high cooling rate ( ⁇ 100 °C / s). This is because if the cooling rate is slow after the end of rolling, the austenite deformed inside the steel sheet can complete the recrystallization process in a short time, at which time the austenite grains grow.
  • the relatively coarse austenite undergoes ferrite transformation during the subsequent cooling process, the ferrite grains formed along the grain boundary of the prior austenite are coarser, usually between 10 and 20 ⁇ m, to increase the strength of the steel sheet. unfavorable.
  • the design idea of the steel plate of the invention is small equiaxed ferrite and martensite structure.
  • the average grain size of the ferrite must be controlled below 10 ⁇ m, which requires the steel plate to be after the finish rolling. It must be cooled quickly to the desired intermediate shutdown temperature. Since the present invention is a low carbon steel, the ferrite transformation has a large driving force and is easily formed. Therefore, the cooling rate after strip rolling should be fast enough ( ⁇ 100 °C / s) to avoid the formation of ferrite during cooling.
  • the cooling temperature of the first stage in the staged cooling process of the present invention needs to be controlled within a temperature range of 600 to 700 ° C. This is because the hot strip rolling line has a fast running speed, and the length of the water cooling stage is limited, and it is impossible to carry out the length. Time is cold.
  • the first stage of the cooling temperature is controlled as much as possible in the optimal temperature range for ferrite precipitation; the main purpose of the second stage water cooling is to form the desired martensite, the martensite transformation is a trimming type transformation, and the transformation speed is fast. Can be completed in an instant, regardless of time.
  • the martensite transformation can be completed as long as the cooling rate reaches the critical cooling rate of the martensite transformation. Therefore, the water cooling rate of the second stage should be controlled between 30 and 50 ° C / s. Excessive cooling rate will cause excessive stress inside the steel plate and poor steel plate shape.
  • the invention can obtain high-strength hot-rolled dual-phase steel with good strength and plasticity through ingenious and reasonable component design and the innovative hot rolling process.
  • the microstructure of the steel plate is fine ferrite and martensite, the volume fraction of ferrite is 20 to 35%, the average grain size of ferrite is 5 to 10 ⁇ m, and the volume fraction of martensite is 65 to 80.
  • the invention adopts a relatively economical component design idea, and at the same time, with the existing hot continuous rolling production line, a high-strength hot-rolled duplex steel having a low yield ratio can be produced.
  • the present invention produces a hot-rolled high-strength dual-phase steel sheet having a yield strength ⁇ 500 MPa, a tensile strength ⁇ 980 MPa, an elongation A 80 ⁇ 12%, and a thickness ⁇ 6 mm, which exhibits excellent strength, ductility and toughness. Matching, excellent forming performance, and low yield ratio can be applied to components such as wheels that require high strength and thinning, and have good application prospects.
  • Figure 1 is a typical metallographic photograph of a steel of Example 1 of the present invention.
  • Figure 2 is a typical metallographic photograph of a steel of Example 2 of the present invention.
  • Figure 3 is a typical metallographic photograph of a steel of Example 3 of the present invention.
  • Figure 4 is a typical metallographic photograph of a steel of Example 4 of the present invention.
  • Figure 5 is a typical metallographic photograph of a steel of Example 5 of the present invention.
  • Table 1 shows the composition of the steel of the embodiment of the present invention
  • Table 2 shows the manufacturing process parameters of the steel of the embodiment of the present invention
  • Table 3 shows the properties of the steel of the example of the present invention.
  • the process flow of the embodiment of the invention is: converter or electric furnace smelting ⁇ vacuum furnace secondary refining ⁇ casting billet or ingot ⁇ steel billet (ingot) heating ⁇ hot rolling+rolling after section cooling ⁇ steel coil, wherein key process parameters are shown in Table 2 .
  • Figures 1 - 5 are typical metallographic photographs of the test steel of Examples 1-5, respectively.
  • the microstructure of the steel sheet of the present invention is fine equiaxed ferrite and martensite (in the figure, the white structure is ferrite, and the dark and gray tissues are martensite).
  • the ferrite grains are mostly distributed in the original austenite grain boundaries, the average grain size is 5-10 ⁇ m, and the martensite equivalent grain size is about 20 ⁇ m.
  • the microstructure and the properties of the steel plate correspond well.
  • the ferrite in the microstructure gives the steel plate a lower yield strength, while the presence of martensite (65 to 80% by volume) gives the steel sheet a high tensile strength.
  • the strength makes the duplex steel of the present invention have the characteristics of easy forming and high strength, and is particularly suitable for fields requiring high strength and thinning of wheels and the like.
  • the present invention can produce a 980 MPa grade ferritic martensitic dual phase steel having a yield strength ⁇ 500 MPa, a tensile strength ⁇ 980 MPa, an elongation A 80 ⁇ 12%, and a lower
  • the yield ratio shows excellent strength, ductility and toughness matching.

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  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
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  • Heat Treatment Of Steel (AREA)

Abstract

一种980MPa级热轧双相钢及其制造方法，该钢化学成分重量百分比为C：0.10～0.20％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.02～0.06％，N≤0.006％，Nb：0.01～0.06％，Ti：0.01～0.05％，其余为Fe和不可避免的杂质，且组成元素同时需满足如下关系：0.05％≤Nb+Ti≤0.10％。钢的微观组织为铁素体+马氏体，铁素体平均晶粒尺寸为5～10μm，马氏体等效晶粒尺寸为15～20μm，其屈服强度≥500MPa，抗拉强度≥980MPa，延伸率A 80≥12％，表现出优异的强度、塑性和韧性匹配，同时具有较低的屈强比，可应用于车轮等需要良好成形性能和高强减薄的部位。

Description

一种980MPa级热轧双相钢及其制造方法 技术领域
本发明属于热轧高强钢领域，具体涉及一种980MPa级热轧双相钢及其制造方法。
背景技术
目前，商用车尤其是重卡汽车的车轮用钢通常采用双相钢进行制造，部分经济型轿车的车轮(包括轮辋和轮辐)为了降低成本也采用钢制车轮。采用高强度双相钢制造车轮可有效降低车轮重量，如与普通Q345钢相比，采用DP600(即抗拉强度600MPa级双相钢)可降低车轮重量约10～15％；而采用抗拉强度780MPa级的DP780双相钢可进一步降低车轮重量约5～10％。目前国内绝大部分车轮厂采用的双相钢主要600MPa以下的低强度双相钢，更高强度的双相钢如DP780则应用得不多。
汽车车轮之所以广泛采用双相钢，主要原因是由于双相钢本身所具有的低屈服强度、高抗拉强度即低屈强比、连续屈服以及良好的加工成形性等，相对较好的扩孔性等。从行业发展的趋势来看，车轮用钢的强度总体仍然是向着更高强度方向发展，而且目前所使用的车轮钢强度普遍不高，在500～600MPa之间。随着国家环保法律法规的日益严格和国家对汽车排放限制措施的实施，除乘用车外，商用车领域的汽车轻量化也已成为行业发展的趋势。针对这一发展趋势，有必要开发更高级别如780MPa和980MPa级热轧双相钢，以满足车轮用钢的未来发展需求。此外，高强度双相钢也可以应用到某些汽车结构件，如汽车大梁、防撞梁等。
目前关于高强度双相钢主要分为两大类：一是冷轧高强双相钢；二是热轧高强双相钢。冷轧高强双相钢的强度目前已达到1180MPa，而热轧高强双相钢强度都没有达到780MPa以上。冷轧和热轧双相钢的这种区别一方面是用户对高强双相钢的需求主要针对冷轧产品，对热轧高强度双相钢的产品需求不强；其次是热轧产线与冷轧产线上存在很大不同，许多在冷 轧产线上能够生产的产品放在热轧产线上难以实现，需要对钢种的成分和工艺重新设计和优化。如前所述，随着行业的不断发展，对热轧高强度双相钢的需求必定会越来越多。
发明内容
本发明的目的在于提供一种980MPa级热轧双相钢及其制造方法，该热轧双相钢的屈服强度≥500MPa，抗拉强度≥980MPa，延伸率A80≥12％，表现出优异的强度、塑性和韧性匹配，可应用在车轮等需要良好成形性能和高强减薄的部位。
为达到上述目的，本发明的技术方案是：
本发明加入较高含量的Si以保证在有限的热轧空冷时间内形成一定数量的铁素体组织且扩大铁素体形成的工艺窗口；Nb和Ti的复合添加主要目的是在精轧阶段最大程度地细化奥氏体晶粒，使得相变之后形成的铁素体更加细小，有利于提高钢板的强度和塑性。本发明通过对组织中铁素体和马氏体含量的精确控制，可获得屈服强度≥500MPa，抗拉强度≥980MPa的高强度铁素体马氏体双相钢。
一种980MPa级热轧双相钢，其化学成分重量百分比为：C：0.10～0.20％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.02～0.06％，N≤0.006％，Nb：0.01～0.06％，Ti：0.01～0.05％，其余为Fe和不可避免的杂质，且上述元素同时需满足如下关系：0.05％≤Nb+Ti≤0.10％。
优选的，所述热轧双相钢的化学成分中：C：0.14～0.18％，以重量百分比计。
优选的，所述热轧双相钢的化学成分中：Si：1.2～1.8％，以重量百分比计。
优选的，所述热轧双相钢的化学成分中：Mn：1.4～1.8％，以重量百分比计。
优选的，所述热轧双相钢的化学成分中：Nb：0.03～0.05％，以重量百分比计。
优选的，所述热轧双相钢的化学成分中：Ti：0.02～0.04％，以重量百 分比计。
进一步，所述热轧双相钢的微观组织为细小的铁素体+马氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；马氏体所占体积分数为65～80％，马氏体铁素体等效晶粒尺寸为15～20μm。
本发明所述热轧双相钢的屈服强度≥500MPa，抗拉强度≥980MPa，延伸率A80≥12％。
在本发明钢的成分设计中：
碳：碳是钢中的基本元素，也是本发明中的重要元素之一。碳扩大奥氏体相区，稳定奥氏体。碳作为钢中的间隙原子，对提高钢的强度起着非常重要的作用，对钢的屈服强度和抗拉强度影响最大。在本发明中，为获得抗拉强度980MPa级的高强双相钢，必须保证碳的含量在0.10％以上；但是碳的含量也不能高于0.2％，否则在热轧两段式冷却过程中难以形成所需数量的铁素体。因此，本发明钢中碳含量须控制在0.1～0.2％之间，优选范围在0.14～0.18％。
硅：硅是钢中的基本元素，同时也是本发明中的重要元素之一。这是因为，要获得抗拉强度达980MPa以上的高强双相钢，一方面要控制铁素体的尺寸和数量，同时还要提高马氏体的强度，这需要在成分设计中适当提高碳和锰的含量。而碳和锰都是扩大奥氏体区、稳定奥氏体的元素，在热轧空冷过程中的很短时间内(通常≤10s)难以形成足够数量的铁素体，这就需要添加较高含量的硅元素。硅的加入可明显促进铁素体形成，扩大铁素体形成的工艺窗口，净化铁素体；同时还可以起到部分强化作用。硅的这种作用必须在其含量达到0.8％以上时才表现出来，但Si的含量也不宜太高，否则轧制后的钢板冲击韧性变差。因此，本发明钢中硅含量控制在0.8～2.0％之间，优选范围在1.2～1.8％之间。
锰：锰也是钢中最基本的元素，同时也是本发明中最重要的元素之一。众所周知，锰是扩大奥氏体相区的重要元素，可以降低钢的临界淬火速度，稳定奥氏体，细化晶粒，推迟奥氏体向珠光体的转变。在本发明中，为保证钢板的强度，锰含量应控制在1.0％以上，锰含量过低，过冷奥氏体不够稳定，空冷时容易转变为珠光体类型的组织；同时，锰的含量也不宜超过2.0％，若超过2.0％，不仅炼钢时容易发生Mn偏析，而且在轧后空冷阶段 不容形成足够数量的铁素体，同时板坯连铸时也容易发生热裂。因此，本发明钢中Mn含量控制在1.0～2.0％，优选范围在1.4～1.8％。
磷：磷是钢中的杂质元素。极易偏聚到晶界上，钢中磷的含量较高(≥0.1％)时，形成Fe2P在晶粒周围析出，降低钢的塑性和韧性，故其含量越低越好，一般控制在0.02％以内较好且不提高炼钢成本。
硫：硫是钢中的杂质元素。钢中的硫通常与锰结合形成MnS夹杂，尤其是当硫和锰的含量均较高时，钢中将形成较多的MnS，而MnS本身具有一定的塑性，在后续轧制过程中MnS沿轧向发生变形，降低钢板的横向拉伸性能。故钢中硫含量越低越好，实际生产时通常控制在0.005％以内。
铝：铝是钢中除C、Si、Mn、P、S五大元素之外另一重要合金元素。铝在本发明中的基本作用主要是在炼钢过程中脱氧。钢中铝的含量一般不低于0.02％；同时，若铝的含量超过0.06％，其细化晶粒的作用反而减弱。根据实际生产过程中铝含量的控制水平，本发明将钢中铝的含量控制在0.02～0.06％即可。
氮：氮在本发明中属于杂质元素，其含量越低越好。氮也是钢中不可避免的元素，通常情况下，若在炼钢过程中不进行特殊控制，钢中氮的残余含量通常≤0.006％。这些固溶或游离的氮元素必须通过形成某种氮化物加以固定，否则游离的氮原子对钢的冲击韧性非常不利，而且在带钢轧制的过程中很容易形成全长性的锯齿裂缺陷。本发明中通过添加钛元素，使之与氮结合形成稳定的TiN从而固定氮原子。因此，本发明钢中氮含量控制在0.006％以内且越低越好。
铌：铌也是本发明中的关键元素之一。由于980MPa及以上级别热连轧双相钢通常需要加入较高的硅以促进轧制空冷段铁素体相的形成，而高硅的添加通常会提高马氏体的脆性。本发明中碳本身的含量虽然≤0.20％，但在析出一定量铁素体之后，铁素体中的碳原子会排出，进入未转变的奥氏体中，使得剩余奥氏体中碳含量富集，最终形成的马氏体实际碳含量较高，马氏体脆性较大，高硅的加入则进一步加剧了这种脆性。因此，高Si型热轧双相钢的低温冲击韧性通常较低，为了尽可能提高高Si型高强双相钢的冲击韧性，合金成分设计中加入微量的铌，通过细化晶粒可有效提高 双相钢的冲击韧性。铌的加入起到两方面作用：一是在高温阶段，固溶的铌对奥氏体晶粒长大起到溶质拖曳作用；二是在精轧阶段，通过铌的碳氮化物钉扎奥氏体晶界，细化奥氏体晶粒，并对最终转变的铁素体和马氏体起到细化作用，提高双相钢冲击韧性。因此，本发明钢中铌含量控制在0.01～0.06％，优选范围在0.03～0.05％。
钛：钛是本发明中的重要元素之一。钛在本发明中主要起两个作用：一是与钢中的杂质元素氮结合形成TiN，起到固氮效果；二是与铌配合，起到最佳的细化奥氏体晶粒的作用。钢中游离的氮原子对钢的冲击韧性非常不利，加入微量钛可将游离的氮固定，但在本发明中钛的含量不宜过多，否则容易形成尺寸较大的TiN，对钢的冲击韧性也是不利的；试验证明，钢中只添加Nb不添加Ti，在连铸生产过程中连铸坯容易发生角裂，添加微量钛则可以有效改善角裂问题；同时，本发明中Nb和Ti的含量只要控制在0.05％≤Nb+Ti≤0.10％成分范围内，即可起到良好的细化晶粒效果且成本较低。因此，本发明钢中钛含量范围控制在0.01～0.05％，优选范围为0.02～0.04％。
氧：氧是炼钢过程中不可避免的元素，对本发明而言，钢中氧的含量通过铝脱氧之后一般都可以达到30ppm以下，对钢板的性能不会造成明显不利影响。因此，将钢中的氧含量控制在30ppm以内即可。
本发明所述980MPa级热轧双相钢的制造方法，其包括如下步骤：
1)冶炼、铸造
按上述化学成分进行冶炼、精炼、铸造成铸坯或铸锭；
2)铸坯或铸锭加热
加热温度1100～1200℃，加热时间1～2小时；
3)热轧+分段冷却+卷取
开轧温度为1030～1150℃，在1000℃以上进行3～5个道次粗轧且累计变形量≥50％；中间坯待温温度为900～950℃，再进行3～5个道次精轧且累计变形量≥70％；终轧温度为800～900℃，终轧结束后以≥100℃/s的冷速将钢板水冷至600～700℃；空冷5～10秒后，再将钢板以30～50℃/s的冷速淬火至≤200℃，卷取后以≤20℃/h的冷速冷却至室温。
本发明的制造工艺设计的理由如下：
在轧制工艺设计上，在粗轧和精轧阶段，轧制过程的节奏应尽量快速完成。在终轧结束后应以高的冷速(≥100℃/s)快速冷却至中间停冷温度。这是因为，轧制结束后若冷却速度较慢，钢板内部形变的奥氏体可在很短的时间内完成再结晶过程，此时奥氏体晶粒发生长大。相对粗大的奥氏体在随后的冷却过程发生铁素体相变时，沿原始奥氏体晶界处形成的铁素体晶粒较为粗大，通常在10～20μm之间，对提高钢板的强度不利。
本发明钢板组织设计思路为细小等轴铁素体和马氏体组织，要达到抗拉强度980MPa级别，铁素体平均晶粒的尺寸必须控制在10μm以下，这就需要钢板在终轧结束后必须快速冷却至所需的中间停冷温度。由于本发明为低碳钢，铁素体相变驱动力大，容易形成。因此，带钢终轧后的冷却速度应足够快(≥100℃/s)，避免在冷却过程中形成铁素体。
本发明分段冷却过程中第一阶段的停冷温度需控制在600～700℃的温度范围内，这是因为热连轧产线带钢运行速度快，水冷阶段的长度有限，不可能进行长时间空冷。第一阶段的停冷温度尽量控制在铁素体析出最佳温度区间；第二阶段水冷的主要目的是形成所需的马氏体，马氏体相变属于切边型转变，转变速度快，可在瞬间完成，与时间基本无关。只要冷却速度达到马氏体转变的临界冷却速度即可完成马氏体相变。因此，第二阶段的水冷速度应控制在30～50℃/s之间。过高的冷却速度会造成钢板内部应力过大，带钢板形不良。
本发明通过巧妙合理的成分设计，同时配合创新性的热轧工艺可获得强度和塑性良好的高强度热轧双相钢。钢板的组织为细小的铁素体和马氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；马氏体所占体积分数为65～80％，马氏体等效晶粒尺寸为15～20μm，在成分设计上，通过理论分析和试验研究，Nb和Ti的总量应满足0.05％≤Nb+Ti≤0.10％，同时配合所要求的轧制工艺，才能获得本发明的低屈强比高强度热轧双相钢，同时具有良好的塑性和良好的冲击韧性。
本发明的有益效果：
(1)本发明采用相对经济的成分设计思路，同时配合现有的热连轧产线就可以生产出具有低屈强比高强度热轧双相钢。
(2)本发明制造出屈服强度≥500MPa，抗拉强度≥980MPa，延伸率 A80≥12％，且厚度≤6mm的热轧高强度双相钢板，该钢板表现出优异的强度、塑性和韧性匹配，以及优异的成形性能，同时具有较低的屈强比，可应用于车轮等需要高强减薄的构件，具有良好的应用前景。
附图说明
图1为本发明实施例1钢的典型金相照片。
图2为本发明实施例2钢的典型金相照片。
图3为本发明实施例3钢的典型金相照片。
图4为本发明实施例4钢的典型金相照片。
图5为本发明实施例5钢的典型金相照片。
具体实施方式
下面结合实施例和附图对本发明做进一步说明。
表1为本发明实施例钢的成分，表2为本发明实施例钢的制造工艺参数，表3为本发明实施例钢的性能。
本发明实施例工艺流程为：转炉或电炉冶炼→真空炉二次精炼→铸坯或铸锭→钢坯(锭)加热→热轧+轧后分段冷却→钢卷，其中关键工艺参数参见表2。
图1-图5分别为实施例1-5试验钢的典型金相照片。从图1-图5可以看出，本发明钢板的显微组织为细小等轴铁素体和马氏体(图中，白色组织为铁素体，深色和灰色组织为马氏体)，铁素体晶粒大多分布在原始奥氏体晶界，平均晶粒尺寸为5～10μm，马氏体等效晶粒尺寸约为20μm。显微组织与钢板的性能可很好的对应，组织中铁素体赋予钢板较低的屈服强度，而马氏体(所占体积分数为65～80％)的存在则赋予钢板具有高的抗拉强度，使得本发明所述双相钢具有易成形且强度高的特点，特别适合车轮等需要高强减薄的领域。
从表3可知，本发明可制造出980MPa级铁素体马氏体双相钢，该双相钢的屈服强度≥500MPa，抗拉强度≥980MPa，延伸率A80≥12％，并具有较低的屈强比，表现出优异的强度、塑性和韧性匹配。
表1         单位：重量百
Figure PCTCN2017088965-appb-000001
表2
Figure PCTCN2017088965-appb-000002
表3
Figure PCTCN2017088965-appb-000003

Claims (11)

  1. 一种980MPa级热轧双相钢，其化学成分重量百分比为：C：0.10～0.20％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.02～0.06％，N≤0.006％，Nb：0.01～0.06％，Ti：0.01～0.05％，其余为Fe和不可避免的杂质，且上述元素同时需满足如下关系：0.05％≤Nb+Ti≤0.10％。
  2. 根据权利要求1所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的化学成分中：C：0.14～0.18％，以重量百分比计。
  3. 根据权利要求1所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的化学成分中：Si：1.2～1.8％，以重量百分比计。
  4. 根据权利要求1所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的化学成分中：Mn：1.4～1.8％，以重量百分比计。
  5. 根据权利要求1所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的化学成分中：Nb：0.03～0.05％，以重量百分比计。
  6. 根据权利要求1所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的化学成分中：Ti：0.02～0.04％，以重量百分比计。
  7. 根据权利要求1-6任一项所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的微观组织为细小的铁素体+马氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；马氏体所占体积分数为65～80％，马氏体等效晶粒尺寸15～20μm。
  8. 根据权利要求1-7任一项所述的980MPa级热轧双相钢，其特征在于，所述热轧双相钢的屈服强度≥500MPa，抗拉强度≥980MPa，延伸率A80≥12％。
  9. 如权利要求1-8任一项所述980MPa级热轧双相钢的制造方法，其包括如下步骤：
    1)冶炼、铸造
    按权利要求1-6任一项化学成分进行冶炼、精炼、连铸成铸坯或铸锭；
    2)铸坯或铸锭加热
    加热温度1100～1200℃，加热时间1～2小时；
    3)热轧+分段冷却+卷取
    开轧温度为1030～1150℃，在1000℃以上进行3～5个道次粗轧且累计变形量≥50％；中间坯待温温度为900～950℃，进行3～5个道次精轧且累计变形量≥70％；终轧温度为800～900℃，终轧结束后以≥100℃/s的冷速将钢板水冷至600～700℃；空冷5～10秒后，再以30～50℃/s的冷速淬火至≤200℃，卷取后以≤20℃/h的冷速冷却至室温。
  10. 根据权利要求9所述的980MPa级热轧双相钢的制造方法，其特征在于，所述热轧双相钢的微观组织为细小的铁素体+马氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；马氏体所占体积分数为65～80％，马氏体等效晶粒尺寸为15～20μm。
  11. 根据权利要求9或10所述的980MPa级热轧双相钢的制造方法，其特征在于，所述热轧双相钢的屈服强度≥500MPa，抗拉强度≥980MPa，A80≥12％。
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