WO2014166323A1 - 一种700MPa级高强度热轧Q&P钢及其制造方法 - Google Patents

一种700MPa级高强度热轧Q&P钢及其制造方法 Download PDF

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WO2014166323A1
WO2014166323A1 PCT/CN2014/073344 CN2014073344W WO2014166323A1 WO 2014166323 A1 WO2014166323 A1 WO 2014166323A1 CN 2014073344 W CN2014073344 W CN 2014073344W WO 2014166323 A1 WO2014166323 A1 WO 2014166323A1
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steel
rolled
strength
hot
mpa
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French (fr)
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王焕荣
李自刚
王巍
张建苏
李建业
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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 KR1020157021206A priority Critical patent/KR101694875B1/ko
Priority to JP2015558341A priority patent/JP6064059B2/ja
Priority to US14/769,647 priority patent/US10023928B2/en
Publication of WO2014166323A1 publication Critical patent/WO2014166323A1/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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/84—Controlled slow cooling
    • 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
    • C21D7/00—Modifying the physical properties of iron or steel by deformation
    • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties 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 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
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • 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
    • C22C38/14—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
    • C21D2211/001—Austenite
    • 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
    • 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 by deformation combined with, or followed by, 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
    • 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
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
    • Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
    • Y02P10/00—Technologies related to metal processing
    • Y02P10/20—Recycling

Definitions

  • the invention belongs to the field of wear-resistant steel, and relates to a 700MPa grade high-strength hot-rolled Q&P steel and a manufacturing method thereof.
  • the yield strength of hot-rolled Q&P steel is ⁇ 700MPa
  • the tensile strength is ⁇ 1300 ⁇ « ⁇
  • the elongation is >10%.
  • Q&P steel Quenching-dispensing steel, Q&P steel, is a research hotspot in the field of high-strength steel for nearly ten years. Its main purpose is to improve the steel's strength while increasing the steel's plasticity, that is, to increase the steel's strong plastic product. At present, Q&P Steel has been recognized as an important new type of steel in the third generation of advanced high strength steel for automotive steel.
  • the main process of Q&P steel is: heating the steel to the complete austenite or part of the austenite zone, and after quenching for a period of time, quenching rapidly to Ms and Mf (Ms and Mf respectively indicate the martensite transformation start temperature and end a temperature between the temperature) to obtain a martensite + retained austenite structure with a certain amount of retained austenite, followed by a certain time for the carbon atoms to pass through at the quenching stop temperature or slightly above the stop cooling temperature
  • the saturated martensite is diffused and enriched in the retained austenite to stabilize the retained austenite and then quenched to room temperature again.
  • Chinese patent CN102226248A discloses a carbon-silicon-manganese hot-rolled Q&P steel, but the alloy composition is not designed for micro-Ti treatment
  • Chinese patent CN101775470A discloses a production process of multi-phase Q&P steel, which is actually a two-step production process. The process of Q&P steel
  • Chinese patent CN101487096A discloses a production of C-Mn-Al based Q&P steel by a two-step heat treatment method, the main feature of which is high elongation, but low strength.
  • the above patent uses a heat treatment method that can be easily controlled by heating in a two-phase zone.
  • the volume fraction of ferrite, but for the hot rolling process the heating temperature is usually in the complete austenite region and the finishing temperature is generally above 780 °C, while the initial precipitation temperature of ferrite is below 700 °C. . Therefore, obtaining a certain amount of ferrite by lowering the finish rolling temperature is difficult to achieve in actual hot rolling production. Summary of the invention
  • the object of the present invention is to provide a 700MPa grade high-strength hot-rolled Q&P steel and a method for producing the same, which have a certain amount of ferrite, martensite and a certain amount of retained austenite structure and have excellent comprehensive performance.
  • the yield strength is ⁇ 700MPa
  • the tensile strength is ⁇ 1300 ⁇ « ⁇
  • the elongation is >10%
  • the alloy cost is greatly reduced. It can be applied to fields requiring easy deformation and moderate wear resistance.
  • the invention adopts a reasonable composition design to suppress the precipitation of cementite by increasing the Si content on the basis of the composition of ordinary C-Mn steel, and the micro-Ti treatment refines the austenite grains, and the A1 content is accelerated to accelerate the austenite process in the air cooling process.
  • the bulk transformation kinetics; at the same time, the hot strip rolling process is combined with the segmental cooling process to obtain the structure containing pro-eutectoid ferrite + martensite + retained austenite.
  • a 700MPa grade high-strength hot-rolled Q&P steel has a chemical composition weight percentage of C: 0.15% ⁇ 0.40%, Si: 1.0% ⁇ 2.0%, Mn: 1.5% ⁇ 3.0%, P ⁇ 0.015%, S ⁇ 0.005%, Al: 0.3% ⁇ 1.0%, N ⁇ 0.006%, Ti: 0.005% ⁇ 0.015%, the rest are Fe and other unavoidable impurities; the yield strength of the 700MPa grade high strength hot rolled Q&P steel ⁇ 700MPa, tensile strength ⁇ 1300 ⁇ « ⁇ , elongation > 10%.
  • Si 1.3-1.7 wt.%; Mn: 1.8-2.5 wt.%; N ⁇ 0.004 wt.%; Ti: 0.008-0.012 wt.%; O ⁇ 30ppm.
  • the action and content control of the chemical composition of the 700 MPa high-strength hot-rolled Q&P steel of the present invention are as follows:
  • Carbon is the most basic element in steel and is one of the most important elements in the 700MPa grade high strength hot rolled Q&P steel of the present invention. 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. In general, the higher the strength of the steel, the lower the elongation. The invention is to ensure that the tensile strength lOOOOMPa is obtained On the high-strength steel plate, the carbon content in the steel is usually not less than 0.15%. The too low carbon content cannot ensure that the carbon is sufficiently saturated from the supersaturated martensite to the retained austenite during the slow cooling of the steel plate after quenching and coiling.
  • the carbon content in the steel should not be too high. If the carbon content is more than 0.4%, although the high strength of the steel can be ensured, the purpose of the present invention is to obtain a certain amount of pro-eutectoid ferrite + martensite + residual In the microstructure, the precipitation of pro-eutectoid ferrite inevitably leads to the enrichment of carbon in the remaining untransformed austenite. The high-carbon martensite elongation obtained after quenching of this part of austenite is too low, so that the final steel plate The elongation is reduced. Therefore, a suitable carbon content in the steel should be controlled at 0.15 ⁇ 0.4wt.% to ensure good strength and good plastic matching of the steel sheet;
  • Silicon is the most basic element in steel and one of the most important elements in the steel of the present invention. Compared with traditional hot-rolled high-strength steels, hot-rolled high-strength Q&P steels are basically designed with high Si. Except for (3, Si, Mn, basically no additional or only a small amount of other alloying elements are added. Si can inhibit the precipitation of cementite in a certain temperature range, but the inhibition of ⁇ carbide is relatively stable and retained austenite. The addition of higher A1 and lanthanum can also inhibit the precipitation of cementite, but the high content of A1 makes the molten steel relatively viscous. It is easy to block the nozzle during continuous casting and reduce the efficiency of pouring steel.
  • the composition design of high Si is still one of the most important component design principles for hot-rolled Q&P steel.
  • the content of Si is generally not less than 1.0 wt.%, otherwise it cannot be suppressed.
  • the cementite is precipitated;
  • the content of Si is generally not more than 2.0 wt.%. Otherwise, the steel plate is prone to thermal cracking during welding, which makes it difficult to apply the steel plate. Therefore, the content of Si in the steel of the present invention is usually controlled at 1.0 to 2.0 wt.%. , preferably in the range of 1.3 to 1.7 wt.%;
  • Manganese is the most basic element in steel and is one of the most important elements in the steel of the present invention. It is well known that Mn is an important element for 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. In order to ensure the strength of the steel plate, the Mn content should generally be controlled above 1.5 wt.%, and the Mn content is too low. When the first stage of the stage cooling is air-cooled, the supercooled austenite is unstable and easily converted into a pearlite type.
  • the content of Mn is generally not more than 3.0 wt.%, Mn segregation is likely to occur during steel making, and hot cracking is likely to occur during slab continuous casting, which is not conducive to the improvement of production efficiency. Therefore, the content of Mn in the steel of the present invention is generally controlled to be 1.5 to 3.0 wt.%, preferably in the range of 1.8 to 2.5 wt.%;
  • Phosphorus is an impurity element in steel. P is easily segregated to the grain boundary, and the content of P in the steel is high. ( ⁇ 0.1wt.%), the formation of Fe 2 P precipitates around the grains, reducing the plasticity and toughness of the steel, so the lower the content, the better, generally controlled within 0.015 wt.% and does not increase the cost of steelmaking ;
  • Sulfur is an impurity element in steel.
  • S in steel usually combines with Mn to form MnS inclusions. Especially when the content of both S and Mn is high, more MnS will be formed in the steel, and MnS itself has a 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 content of S in the steel of the present invention, the better, and the actual production is usually controlled within 0.005 wt.%;
  • Aluminum is one of the most important alloying elements in the steel of the present invention.
  • the basic role of A1 is to deoxidize during the steelmaking process.
  • A1 can also combine with N in steel to form A1N and refine grains.
  • the main purpose of adding more A1 in the steel of the present invention is to accelerate the transformation kinetics of austenite to ferrite in the hollow cooling stage of the staged cooling process, and simultaneously inhibit the cementite with Si. Precipitation, thereby obtaining more metastable retained austenite.
  • the content of A1 in the steel is less than 0.3 wt.%, the ferrite is difficult to be analyzed in a few seconds of air cooling; if the content of A1 in the steel is higher than 1.0 wt.%, the molten steel becomes more viscous, continuous casting The process is prone to block the nozzle and affect production efficiency. Therefore, the content of A1 in the steel of the present invention needs to be controlled within a suitable range, and is controlled at 0.3 to 1.0 wt.%;
  • Nitrogen is an impurity element in the steel of the present invention, and the lower the content, the better. N is also an unavoidable element in steel. Normally, the residual content of N in steel is between 0.002 and 0.004 wt.%. These solid or free N elements can be fixed by binding to acid-soluble A1. In order not to increase the steelmaking cost, the content of niobium may be controlled within 0.006 ⁇ 1.%, preferably in the range of less than 0.004 wt.%; Titanium: The amount of titanium added corresponds to the amount of nitrogen added to the steel.
  • the content of Ti and N in the steel is controlled in a lower range, and a large amount of fine dispersed TiN particles can be formed in the steel during hot rolling; and Ti/N in the steel should be controlled below 3.42 to ensure that all Ti forms TiN. If it is small and has a good height of more than 3.42, it is easy to form relatively coarse TiN particles in the steel, which adversely affects the impact toughness of the steel sheet, and the coarse TiN particles can become a crack source. On the other hand, the content of Ti should not be too low, otherwise the amount of TiN formed is too small to effect the refinement of austenite grains. Therefore, the content of titanium in the steel of the present invention is controlled to be in a suitable range, and the amount of titanium added is 0.005-0.015 wt.%, preferably in the range of 0.008-0.012 wt.%;
  • Oxygen is an inevitable element in the steel making process.
  • the content of 0 in the steel can generally reach 30 ppm or less after deoxidation by A1, and the performance of the steel sheet is not caused. Significant adverse effects. Therefore, the content of 0 in the steel of the present invention can be controlled within 30 ppm.
  • the method for manufacturing the 700MPa grade high-strength hot-rolled Q&P steel of the present invention specifically includes the following steps:
  • the smelting or ingot is formed by converter or electric furnace smelting, secondary refining and vacuum casting according to the following components.
  • the weight percentage of chemical components is: C: 0.15% ⁇ 0.40%, Si: 1.0% ⁇ 2.0%, Mn : 1.5% ⁇ 3.0%, P ⁇ 0.015%, S ⁇ 0.005%, Al: 0.3% ⁇ 1.0%, N ⁇ 0.006%, Ti: 0.005% ⁇ 0.015%, the rest are Fe and other unavoidable impurities;
  • the slab or ingot obtained in step 1) is heated to 1100 ⁇ 1200 °C and kept for l ⁇ 2h, the rolling temperature is 1000 ⁇ 1100°C, multi-pass rolling and cumulative deformation ⁇ 50%, the main purpose is Refine the austenite grains; then carry out 3 ⁇ 5 passes rolling and the cumulative deformation is ⁇ 70% when the intermediate billet temperature drops to 900 ⁇ 950 °C; the hot rolling process is shown in Figure 2; The number of hot rolling passes is 5 ⁇ 7 times;
  • the hot rolled steel is rapidly water cooled to 500-600 °C at a cooling rate of >50 °C /s between 800 °C and 900 °C, then air cooled for 5 ⁇ 10 s, and then continue to cool at >50 °C/s. Cool to a temperature between 100 ⁇ 300 °C (between Ms-Mf) to obtain pro-eutectoid ferrite + martensite + retained austenite structure, and finally slowly cool to room temperature after coiling.
  • the 700MPa grade high strength hot rolled Q&P steel; the post-rolling cooling process is shown in FIG.
  • the multi-pass rolling in step 2) is 5 ⁇ 7 rolling; the slow cooling speed after winding in step 3) is 8 ⁇ 12 °C/h.
  • the volume percentage of the pro-eutectoid ferrite structure is 10-20%, and the volume percentage of the retained austenite structure is more than 5% and less than 10%.
  • the 700MPa grade high-strength hot-rolled Q&P steel of the invention can obtain the steel plate with excellent comprehensive performance by reasonable composition design and the innovative new technology of hot rolling and sectional cooling, that is, the yield strength is ⁇ 700MPa, and the tensile strength is obtained. ⁇ 1300 ⁇ « ⁇ , 700MPa grade high strength hot rolled Q&P steel with elongation >10%.
  • the main purpose of the first stage rapid water cooling is to increase the phase change driving force of the supercooled austenite, and to facilitate the precipitation of sufficient pro-eutectoid ferrite in the subsequent air-cooling stage (10-20) Wt%) to ensure a low yield strength of the steel sheet.
  • the content of carbon and manganese in order to increase the tensile strength of the steel sheet, it is necessary to increase the content of carbon and manganese, and both carbon and manganese are austenite stabilizing elements. Increasing the carbon and manganese content is bound to cause the amount of ferrite precipitated in a limited period of air cooling. Insufficient or even impossible to precipitate.
  • one of the innovations of the present invention greatly increases the aluminum content in the composition design, which is more than ten times that of the general steel.
  • the purpose of substantially increasing the aluminum content is to accelerate the precipitation of ferrite in the air cooling stage at a high carbon and manganese content.
  • the aluminum content should not be too high, otherwise the molten steel will become viscous, and the nozzle will be easily blocked during casting and the alumina inclusions of the steel species will increase. Therefore, the alloy distribution ratio and the hot rolling and cooling processes must be controlled to be well controlled. The faster the water cooling rate in this stage, the better;
  • the second stage quenching and cooling temperature must be controlled within a certain temperature range instead of room temperature. Otherwise, the distribution of carbon atoms cannot be completed and the amount of retained austenite is too low, resulting in a decrease in the elongation of the steel sheet.
  • the commonly used in-line quenching process is directly quenched to room temperature, and another innovation of the present invention is to control the coiling temperature within a certain low temperature range: on the one hand, more retained austenite can be retained (>5 ⁇ 1%), but the retained austenite is unstable at this time. If it is cooled to room temperature, the retained austenite will be transformed into other tissues.
  • adding a certain amount of Si element to the composition design can suppress the carbide in the retained austenite.
  • Precipitation reducing carbon consumption; on the other hand, because the chemical potential of carbon atoms in martensite is higher than the chemical potential in retained austenite, the difference between the chemical potentials of the two is from the martensite to the carbon atoms.
  • the diffusion in the retained austenite provides a driving force, so that the carbon content in the retained austenite is significantly increased, so that the retained austenite can be stably present at room temperature.
  • the heating temperature of the billet is lower than 1100 ° C and the holding time is too short, it is not conducive to the homogenization of the alloying elements; and when the temperature is higher than 1200 ° C, not only the manufacturing cost is increased, but also the heating quality of the billet is Declined. Therefore, the heating temperature of the billet is generally controlled at 1100 ⁇ 1200 °C.
  • the holding time of heating needs to be controlled within a certain range. If the holding time is too short, the diffusion of solute atoms such as Si and Mn is insufficient, and the heating quality of the billet is not guaranteed. The long holding time of the heating makes the austenite grains coarse and the manufacturing cost is increased. Should be controlled between 1 ⁇ 2 hours. The higher the heating temperature, the corresponding holding time can be appropriately shortened.
  • the production process of the invention can be used for manufacturing yield strength ⁇ 700 MPa, tensile strength
  • the alloy cost of the 700 MPa high strength hot rolled Q&P steel sheet of the present invention is greatly reduced. Compared with the traditional low-alloy high-strength steel, no precious metal elements such as Nb, V, Cu, Ni, Mo, etc. are added, which greatly reduces the alloy cost.
  • the production by the hot strip rolling process can further reduce the production cost compared with the production of the thick plate line. Therefore, the production cost of the steel sheet is very low;
  • the 700 MPa high-strength hot-rolled Q&P steel sheet of the present invention has excellent mechanical properties and a reduced user comprehensive use cost. Since the yield strength of the steel sheet is low, the tensile strength is high and the yield ratio is low.
  • One of the biggest benefits of this is that many high-strength steel users can bend the steel plate without modifying existing processing equipment, which saves the cost of retrofitting equipment. It also reduces the loss of abrasive tools and extends the use of abrasive tools. Life expectancy;
  • the steel sheet according to the present invention has the advantages of low cost, low yield ratio and high strength, and is particularly suitable for fields requiring bending forming of a steel sheet and requiring wear resistance.
  • the metastable retained austenite retained in the steel can be transformed into martensite under conditions such as abrasive wear, which further improves the wear resistance of the steel sheet.
  • FIG. 1 is a flow chart of a production process of a 700 MPa high-strength hot-rolled Q&P steel according to the present invention
  • FIG. 2 is a rolling process diagram of a 700 MPa high-strength hot-rolled Q&P steel according to the present invention
  • FIG. 3 is a cooling process diagram of a 700 MPa high-strength hot-rolled Q&P steel after rolling according to the present invention
  • FIG. 4 is a typical metallographic photograph of the test steel of Example 1 of the present invention
  • Figure 5 is a typical metallographic photograph of the 3# test steel of the embodiment of the present invention.
  • Figure 6 is a photograph of a typical metallographic phase of a test steel of Example 5 of the present invention. detailed description
  • the manufacturing method of the 700MPa grade high-strength hot-rolled Q&P steel of the invention has the following production process: converter or electric furnace smelting ⁇ vacuum furnace secondary refining ⁇ casting billet (ingot) ⁇ steel billet (ingot) reheating ⁇ hot rolling+ Sectional cooling process ⁇ steel coil, as shown in Figure 1.
  • Example The manufacture of the 700 MPa high-strength hot-rolled Q&P steel of Examples 1 to 5 specifically includes the following steps:
  • the slab or ingot obtained in step 1) is heated to 1100 ⁇ 1200 °C and kept for l ⁇ 2h, the rolling temperature is 1000 ⁇ 1100 °C, rolling 5 ⁇ 7 times and the cumulative deformation is ⁇ 50%;
  • the intermediate billet temperature is lowered to 900 ⁇ 950 °C, 3 ⁇ 5 passes are rolled and the cumulative deformation is ⁇ 70%;
  • the hot rolling process is shown in Fig. 2; the specific heating and hot rolling process parameters of each embodiment are as follows: Table 2 shows that the thickness of the billet is 120 mm;
  • the hot rolled steel is rapidly water cooled to 500-600 °C at a cooling rate of >50 °C /s between 800 °C and 900 °C, then air cooled for 5 ⁇ 10 s, and then continue to cool at >50 °C / s. Cool to a temperature between 100 ⁇ 300 °C (between Ms-Mf) to obtain a certain amount of ferrite + martensite + - quantitative retained austenite structure, and finally cool slowly to room temperature after coiling (The cooling rate is 8 ⁇ 12 °C / h), the 700MPa grade high-strength hot-rolled Q&P steel of each embodiment is obtained; the post-rolling cooling process is shown in Fig. 3; the specific post-rolling cooling process parameters of each embodiment are shown in the table. 2 is shown.
  • the volume percentages of retained austenite in the steel sheets of Examples 1, 3 and 5 were 5.55%, 6.78% and 8.11%, respectively.
  • the volume fraction of equiaxed pro-eutectoid ferrite is between 10 and 20%.
  • the microstructure of the steel sheet of the present invention is equiaxed pro-eutectoid ferrite + martensite + retained austenite. Due to the presence of retained austenite, the steel sheet undergoes a phase change induced plasticity (TRIP) effect during stretching or wear, thereby improving the wear resistance of the steel sheet.
  • TRIP phase change induced plasticity

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Abstract

一种700MPa级高强度热轧Q&P钢及其制造方法，其化学成分含量为：C：0.15%〜0.40%，Si：1.0%〜2.0%，Mn：1.5%〜3.0%，P≤0.015%，S≤0.005%，Al：0.3%〜1.0%，N≤0.006%，Ti：0.005%〜0.015%，其余为Fe；屈服强度≥700MPa，抗拉强度≥1300MPa，延伸率＞10%。通过合理的成分设计，在普通C-Mn钢成分基础上，通过提高Si含量抑制渗碳体的析出，微Ti处理细化奥氏体晶粒，提高Al含量加快空冷过程的奥氏体转变动力学；同时采用热连轧工艺配合分段冷却工艺，获得含有先共析铁素体+马氏体+残余奥氏体组织。

Description

一种 700MPa级高强度热轧 Q&P钢及其制造方法
技术领域
本发明属于耐磨钢领域， 涉及一种 700MPa级高强度热轧 Q&P钢及 其制造方法， 热轧 Q&P钢的屈服强度≥700MPa, 抗拉强度≥1300^«^, 延 伸率 >10%。
背景技术
淬火-配分钢即 Q&P钢是近十年高强钢领域的研究热点， 其最主要的 目的是在提高钢的强度的同时，提高钢的塑性， 即提高钢的强塑积。 目前， Q&P 钢已经公认为汽车用钢领域第三代先进高强钢中的一类重要新钢 种。
Q&P 钢的主要工艺为： 将钢加热到完全奥氏体区或部分奥氏体区， 均匀化处理一段时间后， 迅速淬火到 Ms和 Mf(Ms和 Mf分别表示马氏体 转变开始温度和结束温度)之间的某一温度以获得具有一定量残余奥氏体 的马氏体 +残余奥氏体组织， 随后在淬火停冷温度或略高于停冷温度下保 温一定时间使碳原子从过饱和的马氏体中向残余奥氏体中扩散富集，从而 稳定残余奥氏体， 然后再次淬火至室温。
Q&P 钢最初的研究和应用主要着眼于汽车行业对高强度高塑性钢材 的需求。 从 Q&P钢的工艺实现过程不难看出， 其工艺路线较为复杂， 钢 板经过第一次淬火之后， 需要快速升温至某一温度并停留一段时间。 这种 两步法 Q&P工艺对于热轧生产过程难以实现， 但是对热轧高强钢的生产 有很好的借鉴意义。 在热轧过程中， 可采用一步法 Q&P工艺即终轧结束 后， 在线淬火至 Ms以下一定温度后卷取。 Q&P钢典型组织为马氏体 +— 定量残余奥氏体， 故具有高强度和良好的塑性。
中国专利 CN102226248A公开了一种碳硅锰热轧 Q&P钢， 但合金成 分设计上没有进行微 Ti处理； 中国专利 CN101775470A公开了一种复相 Q&P钢的生产工艺， 实际上是一种两步法生产 Q&P钢的工艺； 中国专利 CN101487096A公开了一种用两步热处理法生产 C-Mn-Al系 Q&P钢， 其 主要特点是延伸率很高， 但强度较低。
上述专利采用热处理的方法，通过在两相区加热可以较为容易地控制 铁素体的体积分数， 但对于热连轧过程而言， 加热温度通常在完全奥氏体 区且终轧温度一般在 780°C以上， 而铁素体的开始析出温度大多在 700°C 以下。 因此， 通过降低终轧温度来获得一定量的铁素体在热轧实际生产中 难以实现。 发明内容
本发明的目的在于提供一种 700MPa级高强度热轧 Q&P钢及其制造 方法， 该钢种具有一定量的铁素体、 马氏体和一定量的残余奥氏体组织且 综合性能优异， 其屈服强度≥700MPa, 抗拉强度≥1300^«^, 延伸率 >10%, 且合金成本大幅降低， 可应用于要求易变形且中等耐磨的领域。
本发明的设计思路如下：
本发明通过合理的成分设计， 在普通 C-Mn钢的成分基础上， 通过提 高 Si含量抑制渗碳体的析出， 微 Ti处理细化奥氏体晶粒， 提高 A1含量 加快空冷过程的奥氏体转变动力学； 同时采用热连轧工艺配合采用分段冷 却工艺， 获得含有先共析铁素体 +马氏体 +残余奥氏体组织。通过控制三种 不同相的相对含量， 可获得屈服强度 700MPa以上， 抗拉强度在 1300MPa 以上的高强度热轧 Q&P钢。
具体地， 本发明的技术方案是：
一种 700MPa级高强度热轧 Q&P钢，其化学成分的重量百分含量为： C: 0.15%~0.40%, Si: 1.0%~2.0%, Mn: 1.5%~3.0%, P<0.015%, S<0.005%, Al: 0.3%~1.0%, N<0.006%, Ti: 0.005%~0.015%, 其余为 Fe以及其它不 可避免的杂质； 所述 700MPa级高强度热轧 Q&P钢的屈服强度≥700MPa, 抗拉强度≥1300^«^, 延伸率 > 10%。
优选的，所述热轧 Q&P钢的化学成分中， Si: 1.3-1.7 wt.%; Mn: 1.8-2.5 wt.%; N<0.004wt.%; Ti: 0.008-0.012 wt.%; O≤30ppm。
本发明的 700MPa级高强度热轧 Q&P钢的化学成分的作用和含量控 制如下：
碳： 碳是钢中最基本的元素， 同时也是本发明 700MPa级高强度热轧 Q&P 钢中最重要的元素之一。 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。通常情况下， 钢的强度越高， 延伸率越低。 本发明为了保证获得抗拉强度 lOOOMPa 以 上的高强度钢板， 钢中碳的含量通常不低于 0.15%, 过低的碳含量不能保 证钢板在淬火卷取后緩慢冷却过程中碳从过饱和的马氏体向残余奥氏体 中充分扩散， 从而影响残余奥氏体的稳定性。 钢中的含碳量也不宜过高， 若含碳量大于 0.4% , 虽然可以保证钢的高强度， 但由于本发明的目的是 获得一定量先共析铁素体 +马氏体 +残余奥氏体组织，先共析铁素体的析出 必然导致剩余的未转的变奥氏体中富碳，这部分奥氏体在淬火后得到的高 碳马氏体延伸率太低， 从而使得最终钢板的延伸率降低。 因此， 钢中比较 合适的碳含量应控制在 0.15~0.4wt.%可保证钢板具有良好的强度和较好 的塑性匹配；
硅：硅是钢中最基本的元素，同时也是本发明钢中最重要的元素之一。 与传统的热轧高强钢相比， 目前热轧高强 Q&P钢基本都是采用高 Si的成 分设计原则。 除（3、 Si、 Mn外， 基本不添加或仅添加少量其他合金元素。 Si在一定温度范围内可抑制渗碳体的析出，但对 ε碳化物的抑制作用比较 稳定残余奥氏体。 虽然加入较高的 A1和 Ρ也可以抑制渗碳体的析出， 但 A1 含量高使得钢液比较粘稠， 连铸时很容易堵塞水口， 降低浇钢效率； 而 Ρ含量高容易导致晶界脆性， 钢板的沖击韧性很低。 因此， 综合来看， 高 Si 的成分设计仍是目前热轧 Q&P 钢最重要的成分设计原则之一。 Si 的含量一般不低于 1.0 wt.%, 否则不能抑制渗碳体析出； Si的含量一般也 不宜超过 2.0 wt.% , 否则钢板焊接时容易出现热裂， 对钢板的应用造成困 难，故本发明钢中 Si的含量通常控制在 1.0~2.0 wt.% ,优选范围在 1.3~1.7 wt.%;
锰：锰是钢中最基本的元素，同时也是本发明钢中最重要的元素之一。 众所周知， Mn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。 本发明为保证 钢板的强度， Mn含量一般应控制在 1.5 wt.%以上， Mn含量过低， 在分段 冷却的第一阶段空冷时， 过冷奥氏体不稳定， 容易转变为珠光体类型的组 织； 同时， Mn的含量一般也不宜超过 3.0 wt.%, 炼钢时容易发生 Mn偏 析， 同时板坯连铸时易发生热裂， 不利于生产效率的提高。 因此， 本发明 钢中 Mn的含量一般控制在 1.5~3.0 wt.% , 优选范围在 1.8~2.5 wt.%;
磷： 磷是钢中的杂质元素。 P极易偏聚到晶界上， 钢中 P的含量较高 (≥0.1wt.%)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 0.015 wt.% 内较好且不提高炼钢成本；
硫： 硫是钢中的杂质元素。 钢中的 S通常与 Mn结合形成 MnS夹杂， 尤其是档 S和 Mn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。 故本发明钢中 S的含量越低越好， 实际生产时通常控制 在 0.005 wt.% 内；
铝： 铝是本发明钢中最重要的合金元素之一。 A1 的基本作用是在炼 钢过程中进行脱氧。此外， A1还可与钢中的 N结合形成 A1N并细化晶粒。 除上述作用外， 本发明钢中加入较多的 A1的主要目的是加快分段冷却过 程中空冷阶段奥氏体向铁素体的转变动力学过程， 同时与 Si—起共同抑 制渗碳体的析出， 从而获得较多的亚稳态残余奥氏体。 钢中 A1的含量若 低于 0.3 wt.%, 铁素体难以在空冷的几秒钟内充分析出； 若钢中 A1的含 量高于 1.0 wt.%, 钢液变得较为粘稠， 连铸过程容易堵塞水口， 影响生产 效率。因此，本发明钢中 A1的含量需控制在合适的范围内，控制在 0.3~1.0 wt.%;
氮： 氮在本发明钢中属于杂质元素， 其含量越低越好。 N也是钢中不 可避免的元素，通常情况下，钢中 N的残余含量在 0.002~0.004 wt.%之间， 这些固溶或游离的 N元素可以通过与酸溶 A1结合而固定。 为了不提高炼 钢成本，Ν的含量控制在 0.006\¥1.%以内即可，优选范围为小于 0.004wt.%; 钛： 钛的加入量与钢中氮的加入量相对应。 钢中 Ti和 N的含量控制 在较低的范围内， 热轧时可在钢中形成大量细小弥散的 TiN粒子； 同时钢 中 Ti/N需控制在 3.42以下以保证 Ti全部形成 TiN。 细小且具有良好的高 大于 3.42, 则钢中容易形成比较粗大的 TiN粒子，对钢板的沖击韧性造成 不利影响， 粗大的 TiN粒子可成为断裂的裂纹源。 另一方面， Ti的含量 也不能太低， 否则形成的 TiN数量太少， 起不到细化奥氏体晶粒的作用。 因此， 本发明钢中钛的含量要控制在合适的范围， 钛的加入量在 0.005-0.015 wt.%, 优选范围为 0.008-0.012 wt.%;
氧： 氧是炼钢过程中不可避免的元素， 对本发明而言， 钢中 0 的含 量通过 A1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， 将本发明钢中的 0含量控制在 30ppm以内即可。 本发明的 700MPa级高强度热轧 Q&P钢的制造方法， 具体包括如下 步骤：
1 ) 冶炼、 二次精炼、 铸造：
按下述成分采用转炉或电炉冶炼、 真空炉二次精炼、铸造形成铸坯或 铸锭， 化学成分的重量百分含量为： C: 0.15%~0.40%, Si: 1.0%~2.0%, Mn: 1.5%~3.0%, P<0.015%, S<0.005%, Al: 0.3%~1.0%, N<0.006%, Ti: 0.005%~0.015%, 其余为 Fe以及其它不可避免的杂质；
2 ) 力口热、 热库 L:
将步骤 1 )获得的铸坯或铸锭加热到 1100~1200°C并保温 l~2h, 开轧 温度为 1000~1100°C, 多道次轧制且累计变形量≥50%, 主要目的是细化 奥氏体晶粒； 随后待中间坯温度降至 900~950°C时进行 3~5个道次轧制且 累计变形量≥70%； 其热轧工艺如图 2所示； 上述多道次热轧的次数如 5~7 多道次；
3) 分段冷却：
热轧后的轧件在 800~900 °C之间以 >50 °C /s 的冷速快速水冷至 500-600 °C ,然后空冷 5~10s,再继续 >50°C/s的冷速冷却至 100~300°C (即 Ms-Mf之间）之间的某一温度以获得先共析铁素体 +马氏体 +残余奥氏体组 织， 最后卷取后緩慢冷却至室温， 获得所述 700MPa级高强度热轧 Q&P 钢； 其轧后冷却工艺如图 3所示。
较佳的， 步骤 2) 中的多道次轧制为 5~7次轧制； 步骤 3 ) 中的卷取 后緩慢冷却的速度为 8~12°C/h。
上述先共析铁素体 +马氏体 +残余奥氏体组织中，先共析铁素体组织的 体积百分比为 10~20%, 残余奥氏体组织的体积百分比大于 5%且小于 10%。
本发明的 700MPa级高强度热轧 Q&P钢， 通过合理的成分设计， 同 时配合创新性的热轧 +分段冷却的新工艺可获得综合性能优异的钢板， 即 获得屈服强度≥700MPa, 抗拉强度≥1300^«^, 延伸率 >10%的 700MPa级 高强度热轧 Q&P钢。
本发明的分段冷却中，第一阶段快速水冷的主要目的在于提高过冷奥 氏体的相变驱动力，便于在后续空冷阶段析出足够的先共析铁素体（ 10~20 Wt% ) 以保证钢板板较低的屈服强度。 通常为了提高钢板的抗拉强度， 需 要增加碳和锰的含量， 而碳和锰均为奥氏体稳定化元素， 提高碳和锰的含 量势必造成在空冷阶段有限的时间内铁素体析出数量不足甚至无法析出。 因此， 本发明的创新点之一在成分设计上大幅提高了铝的含量， 为一般钢 中铝含量的十倍以上。大幅增加铝含量的目的在于在碳和锰含量较高的情 况下加快空冷阶段铁素体的析出。 但铝含量也不宜过高， 否则容易使钢液 变得粘稠， 浇铸时易堵塞水口且导致钢种氧化铝夹杂增多。 因此， 合金的 成分配比和热轧、 冷却工艺必须控制获得很好地控制， 这一阶段的水冷速 度越快越好；
空冷结束之后第二阶段淬火停冷温度必须控制在一定的温度范围内 而非室温， 否则无法完成碳原子的分配且残余奥氏体数量太低， 导致钢板 延伸率降低。 目前常用的在线淬火工艺都是直接淬火至室温， 而本发明的 又一创新点在于将卷取温度控制在一定的低温范围内：一方面可以保留较 多的残余奥氏体（>5\¥1% ) , 但此时的残余奥氏体不稳定， 若冷却至室温， 残余奥氏体将转变为其他组织， 因此在成分设计上添加一定量的 Si元素 可以抑制残余奥氏体中碳化物的析出， 减少碳的消耗； 另一方面， 由于碳 原子在马氏体中的化学势高于在残余奥氏体中的化学势，二者的化学势之 差为碳原子从马氏体向残余奥氏体中扩散提供了驱动力，使得残余奥氏体 中的碳含量显著增加， 从而使得残余奥氏体在室温下可稳定地存在。 通过 成分配比和冷却工艺的巧妙配合，可获得一定量铁素体 +马氏体 +残余奥氏 体组织的钢板， 从而获得性能优异的 700MPa级高强度热轧 Q&P钢。
另外， 钢坯的加热温度若低于 1100°C以及保温时间过短， 则不利于 合金元素的均匀化； 而当温度高于 1200°C时， 不仅提高了制造成本， 而 且使得钢坯的加热质量有所下降。 因此， 钢坯的加热温度一般控制在 1100~1200 °C比较合适。
同样， 加热的保温时间也需要控制在一定范围内。 保温时间过短， 溶 质原子如 Si、 Mn等的扩散不够充分， 钢坯的加热质量得不到保证； 而加 热的保温时间过长则使得奥氏体晶粒粗大以及提高了制造成本，故保温时 间应控制在 1~2小时之间。 加热温度越高， 相应的保温时间可适当缩短。
本发明的生产工艺可用于制造屈服强度≥700MPa , 抗拉强度
>1300MPa, 且厚度在 3~12mm的高强度热轧 Q&P耐磨钢板， 同时具有良 好的延伸率 (>10%), 该钢板表现出优异的强塑性匹配， 由此带来以下几个 方面的有益效果：
1、 本发明的 700MPa级高强度热轧 Q&P钢板的合金成本大幅降低。 与传统的低合金高强钢相比， 没有添加任何贵重金属元素如 Nb、 V、 Cu、 Ni、 Mo等， 大大降低了合金成本。 采用热连轧工艺生产比用厚板线生产 还可进一步降低生产成本， 因此， 钢板的生产成本很低；
2、 本发明的 700MPa级高强度热轧 Q&P钢板的力学性能优异， 用户 综合使用成本降低。 由于钢板的屈服强度较低， 而抗拉强度高，屈强比低。 这样带来的一个最大好处是，许多高强钢用户无需改造现有加工设备就可 以对钢板进行折弯等工艺， 省却了改造设备的费用； 同时也减少了磨具的 损耗、 延长磨具的使用寿命等；
3、 采用本发明的钢板具有低成本、 低屈强比和高强度优势， 特别适 合于需要对钢板进行折弯成形且要求耐磨的领域。钢中保留的亚稳态残余 奥氏体， 在磨粒磨损等条件下可以转变为马氏体， 进一步提高钢板的耐磨 性。 附图说明
图 1为本发明 700MPa级高强度热轧 Q&P钢生产工艺流程图； 图 2为本发明 700MPa级高强度热轧 Q&P钢轧制工艺图；
图 3为本发明 700MPa级高强度热轧 Q&P钢轧后冷却工艺图； 图 4为本发明实施例 1#试验钢典型金相照片；
图 5为本发明实施例 3#试验钢典型金相照片；
图 6为本发明实施例 5#试验钢典型金相照片。 具体实施方式
下面结合具体实施例对本发明的技术方案进一步详细描述。
本发明的一种 700MPa级高强度热轧 Q&P钢的制造方法， 其生产工 艺流程为： 转炉或电炉冶炼→真空炉二次精炼→铸坯 (锭)→钢坯 (锭)再加 热→热轧 +分段冷却工艺→钢卷， 如图 1所示。 实施例 实施例 1~5的 700MPa级高强度热轧 Q&P钢的制造， 具体包括如下 步骤：
1 ) 冶炼、 二次精炼、 铸造：
按照表 1中各钢的化学成分采用转炉或电炉冶炼、 真空炉二次精炼、 铸造形成铸坯或铸锭；
2 ) 力口热、 热库 L:
将步骤 1 )获得的铸坯或铸锭加热到 1100~1200 °C并保温 l~2h, 开轧 温度为 1000~1100 °C , 5~7次轧制且累计变形量≥50%； 随后待中间坯温度 降至 900~950°C时进行 3~5 个道次轧制且累计变形量≥70%； 其热轧工艺 如图 2所示； 各实施例具体的加热和热轧工艺参数如表 2所示， 钢坯厚度 120mm;
3 ) 分段冷却：
热轧后的轧件在 800~900 °C之间以 >50 °C /s 的冷速快速水冷至 500-600 °C ,然后空冷 5~10s,再继续 >50 °C /s的冷速冷却至 100~300 °C (即 Ms-Mf之间）之间的某一温度以获得一定量铁素体 +马氏体 +—定量的残余 奥氏体组织， 最后卷取后緩慢冷却至室温（冷速为 8~12 °C /h ) , 获得各实 施例的 700MPa级高强度热轧 Q&P钢； 其轧后冷却工艺如图 3所示； 各 实施例具体的轧后冷却工艺参数如表 2所示。
经检测， 实施例 1~5获得的 700MPa级高强度热轧 Q&P钢的力学性 能如表 3所示。 实施例 1、 3和 5的 700MPa级高强度热轧 Q&P钢的典型 金相照片分别如图 4〜图 6所示。
单位： 重量百分比
Figure imgf000009_0001
表 2
Figure imgf000010_0001
表 3 钢板的力学性能
Figure imgf000010_0002
图 4〜图 6的 700MPa级高强度热轧 Q&P钢的典型金相照片上可以清 楚地看出， 钢板的组织主要为等轴状先共析铁素体 +马氏体 +残余奥氏体。
根据 X-射线衍射结果可知， 实施例 1、 3和 5号钢板中残余奥氏体的 体积百分含量分别为 5.55%、 6.78%和 8.11%。 等轴状先共析铁素体的体 积百分含量均在 10~20%之间。 在 500~600°C的温度范围内， 停冷温度越 低， 等轴状先共析铁素体析出量越多。 因此， 本发明钢板的微观组织为等 轴状先共析铁素体 +马氏体 +残余奥氏体。 由于残余奥氏体的存在，使得钢 板在拉伸或磨损过程中发生相变诱导塑性 (TRIP)效应， 从而提高了钢板的 耐磨性。

Claims

权 利 要 求 书
1. 一种 700MPa级高强度热轧 Q&P钢， 其化学成分的重量百分含量为： C: 0.15%~0.40%, Si: 1.0%~2.0%, Mn: 1.5%~3.0%, P<0.015%, S<0.005%, Al: 0.3%~1.0%, N<0.006%, Ti: 0.005%~0.015%, 其余为 Fe以及其它不 可避免的杂质； 所述热轧 Q&P 钢的屈服强度≥700MPa , 抗拉强度 >1300MPa, 延伸率 >10%。
2. 如权利要求 1所述的 700MPa级高强度热轧 Q&P钢， 其特征在于， 所 述热轧 Q&P 钢的化学成分中， Si: 1.3~1.7wt.% ; Mn: 1.8~2.5wt.% ;
N<0.004wt.%; Ti: 0.008~0.012wt.%; O≤30ppm。
3. 如权利要求 1或 2所述的 700MPa级高强度热轧 Q&P钢的制造方法， 具体包括如下步骤：
1 ) 冶炼、 二次精炼、 铸造：
按下述成分采用转炉或电炉冶炼、 真空炉二次精炼、 铸造形成铸 坯或铸锭， 化学成分的重量百分含量为： C: 0.15%~0.40%, Si: 1.0%~2.0% , Mn : 1.5%~3.0% , P<0.015% , S<0.005% , Al： 0.3%~1.0%, N<0.006%, Ti: 0.005%~0.015%, 其余为 Fe以及其 它不可避免的杂质；
2 ) 力口热、 热库 L:
将步骤 1 ) 获得的铸坯或铸锭加热到 1100~1200°C并保温 l~2h, 开轧温度为 1000~1100°C , 多道次轧制且累计变形量≥50%； 随后 待中间坯温度降至 900~950°C时进行 3~5个道次轧制且累计变形 量≥70%；
3 ) 分段冷却：
热轧后的轧件在 800~900 °C之间以 >50 °C /s 的冷速快速水冷至 500-600 °C , 然后空冷 5~10s , 再继续 >50 °C /s 的冷速冷却至 100~300 °C之间的某一温度以获得先共析铁素体 +马氏体 +残余奥 氏体组织， 最后卷取后緩慢冷却至室温， 获得所述 700MPa级高 强度热轧 Q&P钢。
4. 权利要求 3所述的 700MPa级高强度热轧 Q&P钢的制造方法， 其特征 在于， 步骤 2 ) 中的多道次轧制为 5~7次轧制； 步骤 3 ) 中的卷取后緩慢 冷却的速度为 8~12 °C /h。
5. 如权利要求 3所述的 700MPa级高强度热轧 Q&P钢的制造方法， 其特 征在于， 获得的 700MPa级高强度热轧 Q&P钢中， 其先共析铁素体组织 的体积百分比为 10~20%, 残余奥氏体组织的体积百分比大于 5%且小于 10%。
6. 如权利要求 3~5任一所述的 700MPa级高强度热轧 Q&P钢的制造方法， 其特征在于， 获得的 700MPa级高强度热轧 Q&P钢的屈服强度≥700MPa, 抗拉强度≥1300^«^, 延伸率 >10%。
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