WO2017219938A1 - 一种980MPa级热轧铁素体贝氏体双相钢及其制造方法 - Google Patents

一种980MPa级热轧铁素体贝氏体双相钢及其制造方法 Download PDF

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WO2017219938A1
WO2017219938A1 PCT/CN2017/088962 CN2017088962W WO2017219938A1 WO 2017219938 A1 WO2017219938 A1 WO 2017219938A1 CN 2017088962 W CN2017088962 W CN 2017088962W WO 2017219938 A1 WO2017219938 A1 WO 2017219938A1
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bainite
hot
steel
rolled ferritic
mpa
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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 KR1020187038131A priority Critical patent/KR102236344B1/ko
Priority to JP2018566512A priority patent/JP6812465B2/ja
Priority to US16/311,480 priority patent/US11220724B2/en
Publication of WO2017219938A1 publication Critical patent/WO2017219938A1/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
    • 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
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • 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
    • C21D2211/002—Bainite
    • 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

Definitions

  • the invention belongs to the field of hot rolled high strength steel, and particularly relates to a 980 MPa grade hot rolled ferritic bainite 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.
  • Ferritic + bainite duplex steel has a relatively low yield ratio and good hole expandability, plasticity and impact toughness.
  • ferrite + Bayesian Duplex steel has greater application potential than ferritic + martensitic duplex steel.
  • the existing dual-phase steel is mainly ferrite + martensite dual-phase steel, which is mainly cold-rolled ferrite + martensitic duplex steel, hot-rolled ferrite with strength class of 780MPa and above + There are few martensitic duplex steels, and there are fewer high strength ( ⁇ 780 MPa) ferrite + bainite duplex steels.
  • Chinese patent CN101033522A discloses a ferrite bainite dual phase steel whose production is The art is simple, but the composition design contains high aluminum, which is difficult to produce and costly, and its tensile strength is between 700 and 900 MPa.
  • Chinese patent CN102443735A discloses a carbon-manganese ferritic bainite dual phase steel which adopts a segmented cooling process but has a tensile strength of only 450 MPa.
  • Chinese patent CN101603153A discloses a 665MPa grade ferritic bainite dual phase steel, which also adopts a segmented cooling process, but its air cooling time is longer, 12-15 seconds, for thin gauge hot rolled strip. It is difficult to achieve.
  • the object of the present invention is to provide a 980 MPa grade hot-rolled ferritic bainite dual-phase steel and a manufacturing method thereof, wherein the hot-rolled ferritic bainite dual-phase steel has a yield strength of ⁇ 600 MPa and a tensile strength of ⁇ 980 MPa. With an elongation of ⁇ 15%, the duplex steel exhibits excellent strength, plasticity and toughness matching, and can be applied to a part such as a wheel that requires 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 adding higher Al is to be in the air cooling stage after rolling The required amount of ferrite is formed; the main purpose of the composite addition of Nb and Ti is to refine the austenite grains to the greatest extent during the finish rolling stage, so that the ferrite formed after the phase transformation is finer, which is beneficial to the improvement of the steel sheet. Strength and plasticity.
  • the invention can obtain high-strength ferritic bainite dual-phase steel with yield strength ⁇ 600MPa and tensile strength ⁇ 980MPa by precise control of ferrite and bainite content in the structure.
  • a 980 MPa grade hot-rolled ferritic bainite dual-phase steel having a chemical composition weight percentage of C: 0.15 to 0.30%, Si: 0.8 to 2.0%, Mn: 1.0 to 2.0%, P ⁇ 0.02%, S ⁇ 0.005%, O ⁇ 0.003%, Al: 0.5-1.0%, N ⁇ 0.006%, Nb: 0.01-0.06%, Ti: 0.01-0.05%, the rest are Fe and unavoidable impurities, and the above elements are simultaneously satisfied
  • the relationship is as follows: 0.05% ⁇ Nb + Ti ⁇ 0.10%, 2.5 ⁇ Al / C ⁇ 5.0.
  • C 0.20 to 0.25%, by weight percentage.
  • the chemical composition of the hot-rolled ferritic bainite duplex steel Si: 1.2 to 1.8% by weight.
  • Mn 1.4 to 1.8%, In weight percent.
  • Nb 0.03 to 0.05% by weight.
  • the chemical composition of the hot-rolled ferritic bainite duplex steel Ti: 0.02 to 0.04%, by weight percentage.
  • the microstructure of the hot-rolled ferritic bainite dual-phase steel is ferrite + bainite, 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 bainite is 65-80%, and the equivalent grain size of bainite is ⁇ 20 ⁇ m.
  • the hot-rolled ferrite bainite dual-phase steel of the invention has a yield strength of ⁇ 600 MPa, a tensile strength of ⁇ 980 MPa, and an elongation of ⁇ 15%.
  • 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 0.15 and 0.30%, preferably in the range of 0.20 to 0.25%.
  • Silicon is a basic element in steel and is also an important element in the present invention. This is because, in order to obtain a ferritic bainite duplex 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 bainite, which requires design of the composition. Appropriately increase the carbon and manganese content. 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.
  • 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 and can reduce the critical quenching speed of steel. Stabilize austenite, refine grains, and delay the transformation of austenite to pearlite. In the present invention, in order to ensure the strength of the steel sheet, the manganese content should generally 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 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%), Fe2P is formed to precipitate around the grain, which reduces the plasticity and toughness of the steel. Therefore, the lower the content, the better, generally controlled at 0.02. Less than % 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 one of the important alloying elements in the present invention. Since the high-strength ferritic bainite dual-phase steel according to the present invention has higher composition of components, the carbon and manganese contents in the steel are higher than those of other low-strength grade ferritic bainite dual-phase steels. Austenite is more stable, and it is difficult to form ferrite in the staged cooling air cooling stage after rolling, and aluminum is one of the important elements to promote ferrite formation. Therefore, the content of aluminum in the present invention is an order of magnitude higher than that of conventional high-strength steel. The amount of aluminum added to the steel is mainly related to the carbon content, and the amount added should satisfy 2.5 ⁇ Al / C ⁇ 5.0.
  • the aluminum content in the steel of the present invention is controlled to be 0.5 to 1.0%, and the relationship of 2.5 ⁇ Al / C ⁇ 5.0 is satisfied.
  • 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 980 MPa and above hot-rolled ferritic bainite dual-phase steels usually need to be added with 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 bainite. In the present invention, although the content of carbon itself is ⁇ 0.30%, 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. The resulting bainite has a high carbide content and is unfavorable for impact toughness.
  • 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 980MPa grade hot-rolled ferritic bainite 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 3-10 seconds, then 30 ⁇ 50 ° C / s cold water cooling to 350 ⁇ 500 ° C coiling, after coiling to ⁇ 20 ° C / h cooling rate to room temperature.
  • the design idea of the steel plate of the invention is small equiaxed ferrite and bainite 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 required bainite, and the second stage water cooling rate should be controlled at 30 to 50 ° C / s. Excessively high cooling rate will cause the internal stress of the steel plate to be too large, and the strip shape is poor.
  • the coiling temperature can be controlled between 350 and 500 °C.
  • the specific cooling process is shown in Fig. 2.
  • the invention can obtain high-strength hot-rolled ferritic bainite dual-phase steel with good strength and plasticity through ingenious and reasonable component design and the innovative hot rolling process.
  • the structure of the steel plate is fine ferrite and bainite, the volume fraction of ferrite is 20 to 35%, and the average grain size of ferrite 5 ⁇ 10 ⁇ m; bainite accounted for 65-80%, bainite equivalent grain size ⁇ 20 ⁇ m.
  • the total amount of Nb and Ti should meet 0.05 % ⁇ Nb + Ti ⁇ 0.10%; the amount of carbon and aluminum added must satisfy 2.5 ⁇ Al / C ⁇ 5.0, and at the same time with the required rolling process, the low-strength ratio high-strength hot-rolled ferrite of the present invention can be obtained.
  • the invention adopts a relatively economical composition design idea, and at the same time, with the existing hot continuous rolling production line, a high-strength hot-rolled ferrite bainite dual-phase steel with low yield ratio can be produced.
  • the present invention produces a hot-rolled high-strength ferritic bainite dual-phase steel sheet having a yield strength of ⁇ 600 MPa, a tensile strength of ⁇ 980 MPa, an elongation of ⁇ 15%, and a thickness of ⁇ 6 mm, which exhibits excellent strength.
  • plasticity and toughness matching and excellent forming properties while having a low yield ratio, can be applied to components such as wheels that require high strength and thinning, and has a good application prospect.
  • Figure 1 is a schematic view of the heating and rolling process of the present invention.
  • FIG. 2 is a schematic view of a post-rolling cooling process of the present invention.
  • Figure 3 is a typical metallographic photograph of a steel of Example 1 of the present invention.
  • Figure 4 is a typical metallographic photograph of a steel of Example 2 of the present invention.
  • Figure 5 is a typical metallographic photograph of a steel of Example 3 of the present invention.
  • Figure 6 is a typical metallographic photograph of a steel of Example 4 of the present invention.
  • Figure 7 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 3-7 are typical metallographic photographs of the steel of Example 1-5, respectively.
  • the microstructure of the steel sheet of the present invention is fine equiaxed ferrite and bainite (in the figure, the white structure is ferrite and the gray structure is bainite), and the ferrite grains are mostly distributed in the original austenite crystal.
  • the equivalent crystal grain size is 5 to 10 ⁇ m, and the bainite equivalent grain size is about 20 ⁇ m.
  • the microstructure and the properties of the steel plate can be well matched.
  • the ferrite in the microstructure gives the steel plate a lower yield strength, while the presence of more bainite (the volume fraction is 65-80%) gives the steel plate a high
  • the tensile strength makes the ferritic bainite dual-phase steel of the invention have the characteristics of easy forming, good strength, plasticity and toughness matching, and is particularly suitable for fields requiring high strength and thinning of wheels.
  • the present invention can produce a 980 MPa grade ferritic bainite dual phase steel having a yield strength ⁇ 600 MPa, a tensile strength ⁇ 980 MPa, an elongation ⁇ 15%, and a lower yield. Strong ratio, showing excellent strength, plasticity and toughness matching, especially suitable for areas such as wheels that require high strength and thinning.

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  • Heat Treatment Of Steel (AREA)

Abstract

一种980MPa级热轧铁素体贝氏体双相钢及其制造方法，该钢化学成分重量百分比为C：0.15～0.30％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.5～1.0％，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，且同时需满足：0.05％≤Nb+Ti≤0.10％，2.5≤Al/C≤5.0。该钢的微观组织为铁素体+贝氏体，铁素体平均晶粒尺寸为5～10μm，贝氏体等效晶粒尺寸≤20μm，其屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％。

Description

一种980MPa级热轧铁素体贝氏体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级热轧铁素体贝氏体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。然而，铁素体+马氏体型高强双相钢应用于车轮制造的最大缺点是其扩孔性能较差。在相同强度级别水平下，铁素体和马氏体双相钢的扩孔率最低。主要原因是铁素体和马氏体两相的机械性能差异大，加工硬化率高，冲孔周围容易产生微裂纹，在扩孔成形时产生开裂，而相同强度级别的贝氏体或铁素体加贝氏体组织则表现出更优异的扩孔性能。铁素体+贝氏体双相钢具有相对较低的屈强比和良好的扩孔性、塑性和冲击韧性，在超高强度车轮用钢领域(如≥780MPa)，铁素体+贝氏体双相钢具有比铁素体+马氏体双相钢更大的应用潜力。
已有的双相钢主要是铁素体+马氏体型双相钢，这其中又主要是冷轧铁素体+马氏体型双相钢，强度级别在780MPa及以上的热轧铁素体+马氏体型双相钢很少，而高强度(≥780MPa)铁素体+贝氏体型双相钢更少。
中国专利CN101033522A公开了一种铁素体贝氏体双相钢，其生产工 艺简单，但成分设计中含有较高的铝，生产难度较大，成本较高，其抗拉强度在700～900MPa之间。中国专利CN102443735A公开了一种碳锰系铁素体贝氏体双相钢，其采用分段式冷却工艺，但其抗拉强度仅为450MPa。中国专利CN101603153A公开了一种665MPa级铁素体贝氏体双相钢，其也采用分段式冷却工艺，但其空冷时间为较长，为12-15秒，对薄规格热轧带钢而言难以实现。
发明内容
本发明的目的在于提供一种980MPa级热轧铁素体贝氏体双相钢及其制造方法，该热轧铁素体贝氏体双相钢的屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％，该双相钢表现出优异的强度、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；加入较高的Al主要目的是为了在轧后空冷阶段形成所需数量的铁素体；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贝氏体含量的精确控制，可获得屈服强度≥600MPa，抗拉强度≥980MPa的高强度铁素体贝氏体双相钢。
一种980MPa级热轧铁素体贝氏体双相钢，其化学成分重量百分比为：C：0.15～0.30％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.5～1.0％，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％，2.5≤Al/C≤5.0。
优选的，所述热轧铁素体贝氏体双相钢的化学成分中：C：0.20～0.25％，以重量百分比计。
优选的，所述热轧铁素体贝氏体双相钢的化学成分中：Si：1.2～1.8％，以重量百分比计。
优选的，所述热轧铁素体贝氏体双相钢的化学成分中：Mn：1.4～1.8％， 以重量百分比计。
优选的，所述热轧铁素体贝氏体双相钢的化学成分中：Nb：0.03～0.05％，以重量百分比计。
优选的，所述热轧铁素体贝氏体双相钢的化学成分中：Ti：0.02～0.04％，以重量百分比计。
进一步，所述热轧铁素体贝氏体双相钢的微观组织为铁素体+贝氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；贝氏体所占体积分数为65～80％，贝氏体等效晶粒尺寸≤20μm。
本发明所述热轧铁素体贝氏体双相钢的屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％。
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0.15％以上；但是碳的含量也不能高于0.30％，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0.15～0.30％之间，优选范围在0.20～0.25％。
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以上的铁素体贝氏体双相钢，一方面要控制铁素体的尺寸和数量，同时还要提高贝氏体的强度，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，锰含量一般应控制在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：铝是本发明中的重要合金元素之一。由于本发明所涉及高强度铁素体贝氏体双相钢，在成分设计上，其钢中碳和锰的含量相比其他低强度级别的铁素体贝氏体双相钢要更高，奥氏体更稳定，在轧后分段式冷却空冷阶段铁素体形成比较困难，而铝是促进铁素体形成的重要元素之一。因此，本发明中铝的含量较常规高强钢高一个数量级。钢中铝的加入量主要与碳含量相关，其加入量应满足2.5≤Al/C≤5.0。若铝含量偏低，空冷阶段不能形成足够数量的铁素体；铝含量过高，钢水浇铸困难，板坯容易出现表面纵裂等缺陷。因此，本发明钢中铝含量控制在0.5～1.0％，并满足2.5≤Al/C≤5.0的关系式要求。
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及以上级别热连轧铁素体贝氏体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贝氏体的脆性。本发明中碳本身的含量虽然≤0.30％，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，最终形成的贝氏体中碳化物含量高，对冲击韧性不利。为了尽可能提高高Si型铁素体贝氏体双相钢的冲击韧性。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贝氏体起到细化作用，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级热轧铁素体贝氏体双相钢的制造方法，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3～10秒后，再以30～50℃/s的冷速水冷至350～500℃卷取，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：
本发明轧制工艺示意图参见图1。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贝氏体组织，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贝氏体，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，卷取温度控制在350～500℃之间即可，具体冷却工艺示意图如图2所示。
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铁素体贝氏体双相钢。钢板的组织为细小的铁素体和贝氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸 为5～10μm；贝氏体所占体积分数为65～80％，贝氏体等效晶粒尺寸≤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％；碳和铝的添加量须满足2.5≤Al/C≤5.0，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铁素体贝氏体双相钢，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铁素体贝氏体双相钢。
(2)本发明制造出屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％，且厚度≤6mm的热轧高强度铁素体贝氏体双相钢板，该钢板表现出优异的强度、塑性和韧性匹配和优异的成形性能，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为本发明加热和轧制工艺示意图。
图2为本发明轧后冷却工艺示意图。
图3为本发明实施例1钢的典型金相照片。
图4为本发明实施例2钢的典型金相照片。
图5为本发明实施例3钢的典型金相照片。
图6为本发明实施例4钢的典型金相照片。
图7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3-图7分别为实施例1-5钢的典型金相照片。从图3-图7可以看出， 本发明钢板的显微组织为细小等轴铁素体和贝氏体(图中，白色组织为铁素体、灰色组织为贝氏体)，铁素体晶粒大多分布在原始奥氏体晶界，等效晶粒尺寸为5～10μm，贝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，而较多贝氏体(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，使得本发明所述铁素体贝氏体双相钢具有易成形、强度、塑性和韧性匹配良好等特点，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级铁素体贝氏体双相钢，该双相钢的屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％，并具有较低的屈强比，表现出优异的强度、塑性和韧性匹配，特别适合车轮等需要高强减薄的领域。
Figure PCTCN2017088962-appb-000001
表3
Figure PCTCN2017088962-appb-000002

Claims (11)

  1. 一种980MPa级热轧铁素体贝氏体双相钢，其化学成分重量百分比为：C：0.15～0.30％，Si：0.8～2.0％，Mn：1.0～2.0％，P≤0.02％，S≤0.005％，O≤0.003％，Al：0.5～1.0％，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％，2.5≤Al/C≤5.0。
  2. 根据权利要求1所述的980MPa级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体双相钢的化学成分中：C：0.20～0.25％，以重量百分比计。
  3. 根据权利要求1所述的980MPa级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体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级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体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级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体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级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体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级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体双相钢的微观组织为铁素体+贝氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；贝氏体所占体积分数为65～80％，贝氏体等效晶粒尺寸≤20μm。
  8. 根据权利要求1-7任一项所述的980MPa级热轧铁素体贝氏体双相钢，其特征在于，所述热轧铁素体贝氏体双相钢的屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％。
  9. 如权利要求1-8任一项所述980MPa级热轧铁素体贝氏体双相钢的制造方法，其包括如下步骤：
    1)冶炼、铸造
    按权利要求1-6任一项所述化学成分进行冶炼、精炼、铸造坯或铸锭；
    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3～10秒后，再以30～50℃/s的冷速水冷至350～500℃卷取，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级热轧铁素体贝氏体双相钢的制造方法，其特征在于，所述热轧铁素体贝氏体双相钢的微观组织为铁素体+贝氏体，铁素体所占体积分数为20～35％，铁素体平均晶粒尺寸为5～10μm；贝氏体所占体积分数为65～80％，贝氏体等效晶粒尺寸≤20μm。
  11. 根据权利要求9或10所述的980MPa级热轧铁素体贝氏体双相钢的制造方法，其特征在于，所述热轧铁素体贝氏体双相钢的屈服强度≥600MPa，抗拉强度≥980MPa，延伸率≥15％。
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