WO2020248459A1 - 一种高强度钢的热处理方法和由此获得的产品 - Google Patents

一种高强度钢的热处理方法和由此获得的产品 Download PDF

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WO2020248459A1
WO2020248459A1 PCT/CN2019/111796 CN2019111796W WO2020248459A1 WO 2020248459 A1 WO2020248459 A1 WO 2020248459A1 CN 2019111796 W CN2019111796 W CN 2019111796W WO 2020248459 A1 WO2020248459 A1 WO 2020248459A1
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steel
less
strength
carbide particles
strength steel
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PCT/CN2019/111796
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English (en)
French (fr)
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易红亮
常智渊
刘钊源
杨达朋
熊小川
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育材堂（苏州）材料科技有限公司
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Priority to JP2021573256A priority Critical patent/JP7190216B2/ja
Priority to BR112021025011A priority patent/BR112021025011A2/pt
Priority to KR1020227000319A priority patent/KR20220019264A/ko
Priority to AU2019450666A priority patent/AU2019450666A1/en
Priority to EP19932761.0A priority patent/EP3981894A4/en
Priority to US17/617,681 priority patent/US20220251673A1/en
Priority to CA3142958A priority patent/CA3142958A1/en
Publication of WO2020248459A1 publication Critical patent/WO2020248459A1/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
    • C21D1/26—Methods of annealing
    • 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
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • 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
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • 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
    • C21D1/26—Methods of annealing
    • C21D1/28—Normalising
    • C—CHEMISTRY; METALLURGY
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/78—Combined heat-treatments not provided for above
    • C—CHEMISTRY; METALLURGY
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    • 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
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
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    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/02—Hardening by precipitation
    • 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
    • C21D9/0075—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for rods of limited length
    • 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
    • C21D9/02—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for springs
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
    • F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
    • F16—ENGINEERING ELEMENTS AND UNITS; GENERAL MEASURES FOR PRODUCING AND MAINTAINING EFFECTIVE FUNCTIONING OF MACHINES OR INSTALLATIONS; THERMAL INSULATION IN GENERAL
    • F16F—SPRINGS; SHOCK-ABSORBERS; MEANS FOR DAMPING VIBRATION
    • F16F1/00—Springs
    • F16F1/02—Springs made of steel or other material having low internal friction; Wound, torsion, leaf, cup, ring or the like springs, the material of the spring not being relevant
    • F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
    • F16—ENGINEERING ELEMENTS AND UNITS; GENERAL MEASURES FOR PRODUCING AND MAINTAINING EFFECTIVE FUNCTIONING OF MACHINES OR INSTALLATIONS; THERMAL INSULATION IN GENERAL
    • F16F—SPRINGS; SHOCK-ABSORBERS; MEANS FOR DAMPING VIBRATION
    • F16F1/00—Springs
    • F16F1/02—Springs made of steel or other material having low internal friction; Wound, torsion, leaf, cup, ring or the like springs, the material of the spring not being relevant
    • F16F1/021—Springs made of steel or other material having low internal friction; Wound, torsion, leaf, cup, ring or the like springs, the material of the spring not being relevant characterised by their composition, e.g. comprising materials providing for particular spring properties
    • 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
    • C21D2211/004—Dispersions; Precipitations
    • 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 present invention relates to the heat treatment of high-strength steel. More specifically, the present invention relates to a heat treatment method for high-strength steel that enables the processed steel to have high strength, high ductility, and high toughness at the same time.
  • the processed steel is particularly suitable for preparing spring components, such as springs for vehicle suspensions. member.
  • Spring members for vehicle suspensions include, for example, leaf springs, stabilizer bars, round springs, and the like.
  • Leaf springs are abbreviated as leaf springs and are often installed between the frame and the axle.
  • the stabilizer bar is a torsion bar spring.
  • the stabilizer bar uses the elastic force of the shaft to prevent the wheels from lifting, to prevent excessive lateral roll of the body, and try to keep the body balanced.
  • the spring member for vehicle suspension repeatedly bears the stress load.
  • the steel used for preparing the spring member for vehicle suspension needs to have high strength.
  • Chinese patent application CN108239726A provides a high-strength spring steel with excellent hydrogen embrittlement resistance and its manufacturing method.
  • the chemical composition is calculated by weight percentage, and the steel includes: C 0.45-0.60%, Si 1.40-1.80%, Mn 0.30-0.80%, Cr 0.20-0.70%, Mo 0.05-0.15%, V 0.05-0.20%, Nb 0.010-0.030%, N ⁇ 0.006%, P ⁇ 0.015%, S ⁇ 0.015%.
  • the corresponding heat treatment process is as follows: heat the steel to 880-1000°C for 10s-30min, then cool it at an average cooling rate of 10°C/s or more, and then heat the steel to 380-460°C for 10s-40min, Then cool to below 60°C at an average cooling rate of 10°C/s or more.
  • CN108239726A believes that by adding Nb, V, Mo and other strong carbide forming elements, a certain amount of (V, Mo) C or (Nb, V, Mo) C particles with an average grain diameter of 10-60 nm is precipitated in the steel. And the original austenite grains are also refined to more than 10 levels.
  • the carbide particles have high interfacial activation energy and can play a role of non-diffusible hydrogen trapping that does not diffuse due to external stress.
  • the small pristine austenite grains and a sufficient amount of nano-sized carbide particles ensure that the spring steel has a tensile strength above 1900MPa and excellent hydrogen embrittlement resistance.
  • Chinese patent application CN106399837A provides a steel for hot stamping and a hot forming process.
  • the chemical composition is calculated by weight percentage, and the steel contains: C 0.27-0.40%, Si 0-0.80%, Mn 0.20-3.0%, V 0.10-0.4%, Si 0-0.8%, Al 0-0.5%, Cr 0 -2%, Ti 0-0.15%, Nb 0-0.15%, B 0-0.004% and Mo, Ni, Cu with a total content of less than 2%.
  • the nano-scale VC particles and/or the composite carbide particles of V and Ti, Nb are controlled to precipitate during the hot forming process to achieve precipitation strengthening and crystallization.
  • the properties of the steel are further optimized.
  • the resulting steel has a yield strength of 1350-1800MPa, a tensile strength of 1700-2150MPa, and an elongation of 7-10%.
  • the reduction of area is often used as an important index to comprehensively characterize the ductility and toughness of steel.
  • the steel needs to have both high ductility and high toughness.
  • the ductility and toughness of the currently disclosed steel materials are not sufficient to provide the reduction of area performance suitable for preparing spring members for vehicle suspensions. Therefore, there is still a need for a heat treatment method for high-strength steel, which can make the treated steel have high strength, high ductility and high toughness at the same time, especially with high reduction of area, so that it is particularly suitable for preparing vehicle suspensions.
  • Spring member
  • the present invention provides a heat treatment method for obtaining high-strength steel with high reduction of area, which solves the above-mentioned problems.
  • the present invention provides a heat treatment method for high-strength steel, wherein the high-strength steel comprises 0.30-0.45% C, 1.0% or less Si, 0.20-2.5% Mn, by weight percentage, 0.20-2.0% Cr, 0.15-0.50% Mo, 0.10-0.40% V, 0.2% or less Ti, 0.2% or less Nb, and the rest are Fe and other alloying elements and impurities.
  • the alloy composition makes the formula (1 )’S Eq(Mn) is not less than 1.82,
  • the method includes the following steps:
  • Austenitizing step heating the high-strength steel to about 20°C above the austenitizing critical temperature (Ac3) (hereinafter referred to as Ac3+20°C) to about 950°C, and holding for about 1-300 minutes;
  • the high-strength steel is initially cooled to about 10°C below the ferrite precipitation start temperature (Ar3) (hereinafter referred to as Ar3-10°C) to about 870°C, Keep holding for about 5-300 minutes, and then further cool to below about 100°C, wherein the average cooling rate of the further cooling is not less than about 1°C/s; and
  • Ar3-10°C ferrite precipitation start temperature
  • Tempering step After the carbide precipitation step, the high-strength steel is heated again to about 120-280° C., and the temperature is maintained for about 5-360 min.
  • the present invention provides a steel obtained by the above heat treatment method, wherein, in terms of area, the microstructure of the steel includes: about 90% or more of martensite, and about 3% or less of ferrite. , Less than or equal to about 5% of retained austenite, less than or equal to about 10% of bainite.
  • the steel contains about 0.1-0.5% by weight of carbide particles, wherein the carbide particles comprise composite carbide particles of V and Mo, and the average particle size of the carbide particles is about 1-30 nm, and
  • the yield strength of the steel is greater than or equal to about 1400 MPa, the tensile strength is greater than or equal to about 1800 MPa, and the reduction of area is greater than or equal to about 38%.
  • the present invention provides a spring member for a vehicle suspension made of the aforementioned steel.
  • Fig. 1 is a schematic diagram of temperature-time used in an embodiment of the heat treatment method of high-strength steel of the present invention
  • Figure 2 is a metallographic photograph of an embodiment of the steel obtained by the heat treatment method of the present invention.
  • Fig. 3 is a transmission electron microscope photograph of an embodiment of the steel obtained by the heat treatment method of the present invention.
  • Fig. 4 is the result of the carbide chemical composition of an embodiment of the steel obtained by the heat treatment method of the present invention.
  • the present invention relates to a heat treatment method of high-strength steel, wherein the high-strength steel comprises 0.30-0.45% C, 1.0% or less Si, 0.20-2.5% Mn, in terms of weight percentage, 0.20-2.0% Cr, 0.15-0.50% Mo, 0.10-0.40% V, 0.20% or less Ti, 0.2% or less Nb, the rest are Fe and other alloying elements and impurities.
  • the alloy composition makes the formula (1 )’S Eq(Mn) is not less than 1.82,
  • the method includes the following steps:
  • Austenitizing step heating the high-strength steel to about Ac3+20°C to about 950°C, holding for about 1-300min; preferably heating the high-strength steel to about Ac3+30°C to 910°C, Keep warm for about 1-30min;
  • the high-strength steel is initially cooled to about Ar3-10°C to about 870°C, kept for about 5-300 minutes, and then further cooled to below about 100°C, wherein The average cooling rate of the further cooling is not less than about 1°C/s; preferably, the high-strength steel is initially cooled to about Ar3+10°C to 850°C and held for about 5-30 minutes; and
  • Tempering step After the carbide precipitation step, the high-strength steel is heated again to about 120-280°C for about 5-360 minutes; preferably, the high-strength steel is heated again to about 160-230°C for insulation About 10-60min.
  • the high-strength steel preferably contains: 0.32-0.42% C, 0.8% or less Si, 0.2-1.5% Cr, 0.2-0.4% Mo, 0.12-0.3% V, and the rest is Fe And other alloying elements and impurities, where the alloy composition makes the Eq(Mn) of formula (1) not less than 1.82.
  • composition of the high-strength steel used in the present invention is described in detail as follows.
  • C is the most effective solid solution strengthening element in steel.
  • the content of C In order to ensure that the tensile strength of steel is above 1800MPa, the content of C must be greater than or equal to about 0.30%. However, if the C content exceeds 0.45%, the high-carbon martensite formed has poor ductility and toughness, and the hydrogen embrittlement resistance is significantly reduced. Therefore, the C content of the high-strength steel used in the present invention is between about 0.30-0.45%; preferably between about 0.32-0.42%.
  • Si about 1.0% or less
  • Si is a deoxidizer during the smelting of steel, and solid solution in the ferrite matrix has the effect of strengthening the strength of the base material.
  • excessive Si is not only harmful to the toughness of the steel, but also forms serious surface oxidation and decarburization during heat treatment.
  • the thickness of the decarburized layer is one of the key control parameters for the fatigue performance of the spring components for vehicle suspension. Therefore, the Si content of the high-strength steel used in the present invention is 1.0% or less, preferably 0.8% or less.
  • Mn is an element that improves the hardenability of steel and ensures strength.
  • the Mn content is less than about 0.20%, the hardenability of the steel material is insufficient, and it is difficult to obtain high strength.
  • the Mn content is too high, the ductility and toughness of the steel will be significantly reduced. Therefore, the upper limit of the Mn content of the high-strength steel used in the present invention is about 2.5%.
  • the Cr content of the high-strength steel used in the present invention ranges from about 0.20-2.0%, preferably 0.2-1.5%.
  • Mo is a strong carbide forming element and has a greater affinity for carbon atoms. It can prevent the diffusion of carbon atoms and reduce the diffusion coefficient of carbon elements, thereby effectively inhibiting the surface decarburization of steel.
  • the thickness of the decarburized layer is one of the key control parameters for the fatigue performance of the spring components for vehicle suspension.
  • the addition of Mo also improves the hardenability of the steel, and it takes advantage of the precipitation of composite nano-carbides of Mo and V during the heat treatment process.
  • the precipitation of composite carbides is more conducive to the dispersion of carbides and to obtain finer carbides. This not only ensures the ultra-high strength of the steel, but also makes the material have good performance in terms of reduction of area.
  • the content of Mo added in the steel of the present invention is not less than about 0.15%. However, if the Mo content is higher than about 0.50%, the production cost will increase significantly. Therefore, the Mo content is about 0.15-0.50%, and the Mo content is preferably greater than about 0.20 to about 0.40%. When the Mo content is greater than about 0.20 to about 0.40%, the surface decarburization problem of the spring member can be effectively suppressed or alleviated, thereby providing the spring member with good fatigue resistance, and the addition of Mo in this range can also ensure that Mo The composite carbide with V is dispersedly distributed and small in size, so as to provide a good reduction of area for the spring member.
  • V about 0.10-0.40%
  • V forms composite carbides, which play a role in precipitation strengthening and original austenite grain refinement. If the V content is less than about 0.10%, sufficient carbides cannot be formed, and the above effect is not significant. If the V content is higher than about 0.40%, it will lead to an increase in production costs and coarse carbides to decrease the reduction of area. Therefore, the V content is preferably about 0.10-0.40%, preferably 0.12-0.30%.
  • Ti and Nb form carbonitrides in steel, which have an impact on the improvement of strength and the refinement of grains. Because Ti and Nb are the strongest carbide forming elements, when the content exceeds about 0.20%, a large amount of carbonitrides will be precipitated at high temperatures. This causes its size to be coarse, resulting in a decrease in the reduction of area. If the heat treatment is used to control Ti or Nb to precipitate carbonitrides at high temperatures to minimize the precipitation of carbonitrides at high temperatures, it is beneficial to the composite precipitation with V and Mo and can further refine the size of carbonitrides. However, such process control is more complicated. Therefore, the Ti content is about 0.20% or less, preferably 0.05% or less; the Nb content is about 0.20% or less, preferably about 0.05% or less.
  • both Ti and Nb have the effect of forming carbonitrides in the steel and improving the overall performance of the steel.
  • the inventor also found that a synergistic effect can be achieved by compounding the two. Therefore, the total total amount of Ti+Nb is about 0.20% or less. If the total amount of Ti+Nb exceeds about 0.20%, a large amount of carbonitrides will precipitate at a high temperature, resulting in a coarse size and a decrease in the reduction of area. Therefore, the total total amount of Ti+Nb is about 0.20% or less, preferably about 0.08% or less.
  • P and S will segregate at the grain boundary, which leads to a decrease in the reduction of area of steel. Therefore, it is desirable that these elements are as few as possible, for example, the contents of P and S are both less than or equal to about 0.025%.
  • Eq(Mn) characterizes the hardenability of steel.
  • the relationship between Eq(Mn) and the critical cooling rate Rc of martensite satisfies the following formula (2).
  • the balance component of the present invention is iron (Fe).
  • Fe iron
  • impurities from raw materials or the surrounding environment will inevitably be mixed in during the conventional manufacturing process. Therefore, it cannot be ruled out that these impurities are mixed in. These impurities are known to those of ordinary skill in the art.
  • the heat treatment method of the present invention includes the following steps for the above-mentioned high-strength steel: 1) austenitizing step; 2) carbide precipitation step; and 3) tempering step.
  • high-strength steel undergoes an austenitizing step.
  • the austenitizing step is carried out by heating the high-strength steel to about Ac3+20°C to about 950°C and holding the temperature for about 1-300 minutes.
  • the heating temperature is lower than Ac3+20°C or the holding time is less than 1 min, there may be residual ferrite and pearlite that are not solid-solved, resulting in inhomogeneous alloying elements.
  • the heating temperature is higher than 950°C or the holding time exceeds 60min, it may cause serious oxidation and decarburization of the steel surface, and the austenite grains are coarse.
  • the process conditions of the austenitizing step are set to hold at Ac3+20°C-950°C for 1-60 min, preferably at Ac3+30°C-910°C for 1-30 min.
  • the austenitizing step is achieved by holding for 5-310 minutes in a heating furnace with a furnace temperature of Ac3+40°C to about 970°C, preferably for 5-40 minutes in a heating furnace with a furnace temperature of Ac3+50°C to about 930°C achieve.
  • the aforementioned austenitizing step can also be completed by applying induction heating or a combination of induction heating and heating furnace heating.
  • the high-strength steel undergoes a carbide precipitation step.
  • the carbide precipitation step is performed as follows: the high-strength steel is initially cooled to about Ar3-10°C to about 870°C, and the temperature is maintained for about 5-300 minutes; preferably, the high-strength steel is initially cooled to about Ar3+10°C to 850°C Keep the temperature for 5-30 minutes, and then further cool to below about 100°C, wherein the average cooling rate of the further cooling is greater than or equal to about 1°C/s, preferably greater than or equal to about 5°C/s.
  • the average cooling rate of the further cooling may be less than or equal to about 100°C/s, preferably less than or equal to about 50°C/s, more preferably less than or equal to about 20°C/s.
  • the initial cooling can be achieved in a furnace with a furnace temperature of about Ar3-20°C to about 870°C for about 5-300 minutes, preferably in a furnace with a furnace temperature of about Ar3°C to about 850°C for about 5-30 minutes achieve.
  • the further cooling may be performed by means including oil quenching, salt water quenching, and the like. When the holding temperature of the initial cooling of the carbide precipitation step is lower than Ar3-10°C, more ferrite may be formed, which will be harmful to the strength and fatigue properties of the steel.
  • the process conditions of the carbide precipitation step are set to be maintained at about Ar3-10°C-870°C for about 5-300 minutes, and then further cooled to below about 100°C.
  • the average cooling rate of the cooling process is set to be greater than or equal to about 1°C/s.
  • the present invention there are a large number of nano carbide particles in the initial structure of the high-strength steel processed by the present invention. Therefore, after the austenitizing step, a large number of nano carbide particles remain undissolved in the structure of the high-strength steel, which is beneficial for controlling the precipitation amount and size of carbide particles in the carbide precipitation step.
  • the high-strength steel undergoes a tempering step.
  • the tempering step is performed by reheating the high-strength steel to about 120-280° C. and holding it for about 5-360 min.
  • the tempering temperature is lower than about 120°C or the holding time is less than about 5 minutes, the tempering effect of martensite is not sufficient, the internal stress caused by the phase transformation of martensite cannot be fully released, and the reduction of area performance cannot be obtained.
  • the tempering temperature is higher than about 280°C or the holding time exceeds about 360min, it will cause a large amount of precipitation of Fe-C carbides, which will result in a significant decrease in the strength of the steel and a significant tendency of carbide coarsening.
  • the manufacturing process of automobile components includes a step of baking after painting.
  • the baking is performed by keeping the temperature at about 150-230°C for about 10-60 minutes.
  • the baking step can function as the above-mentioned tempering step, so no additional tempering step is required.
  • the microstructure of the steel material includes: about 90% or more martensite, about 3% or less ferrite, and about 5% or less.
  • Retained austenite of less than or equal to about 10% of bainite preferably includes greater than or equal to about 97% of martensite, and the sum of retained austenite, ferrite and bainite is less than or equal to about 2.5%;
  • the steel contains about 0.1-0.5% by weight of carbide particles, wherein the carbide particles comprise composite carbide particles of V and Mo, and the average particle size of the carbide particles is about 1-30 nm, and
  • the yield strength of the steel is greater than or equal to about 1400 MPa, the tensile strength is greater than or equal to about 1800 MPa, and the reduction of area is greater than or equal to about 38%; preferably, the yield strength of the steel is greater than or equal to about 1550 MPa, and the tensile strength is greater than or equal to about 1900 MPa, And the reduction of area is greater than or equal to about 45%.
  • the steel material obtained by the heat treatment method of the present invention is described in detail as follows.
  • the microstructure includes about 90% or more martensite.
  • Martensite is the microstructure required to obtain high strength.
  • the area percentage of martensite is less than about 90%, it means that there are too many ferrite and retained austenite that contribute little to the increase in strength, making it difficult to achieve high tensile strength.
  • the microstructure includes about 90% or more of martensite, thereby ensuring the strength of the steel.
  • the area percentage of the martensite is preferably about 97% or more, and may be about 99% or more.
  • the microstructure includes about 10% or less of bainite.
  • Bainite has a lower hardness than martensite.
  • the presence of bainite in the steel will reduce the strength of the steel. Therefore, the bainite content should not exceed 10%.
  • the bainite content is about 3% or less, and may be about 0%.
  • the microstructure includes about 3% or less of ferrite.
  • Ferrite is a soft phase. When it and martensite coexist in steel, the two form a larger hardness difference, which will significantly reduce the strength of the steel. Therefore, the generation of ferrite should be avoided as much as possible, and the ferrite content is preferably less than or equal to about 1%, and may be about 0%. .
  • the microstructure includes about 5% or less of retained austenite.
  • Retained austenite can improve the ductility of steel and help to improve the hydrogen embrittlement resistance of steel. Therefore, the steel material of the present invention may contain a certain amount of retained austenite. But the retained austenite will reduce the strength of the steel, so it should not be too much. Excessive retained austenite will form high-carbon martensite when the steel is plastically deformed, which is harmful to the toughness of the steel.
  • the retained austenite is preferably 3% or less, and more preferably 1% or less.
  • the steel obtained in the present invention contains about 0.1-0.5% by weight of carbide particles, wherein the carbide particles comprise composite carbide particles of V and Mo, and the average particle size of the particles is about 1-30 nm.
  • carbide particles will precipitate on the austenite grain boundaries, which pin the austenite grains, thereby inhibiting the growth of the austenite grains.
  • carbide particles are also precipitated in the austenite grains.
  • the second phase particles are more uniform and finer, which can play a role of precipitation strengthening and increase the strength of the steel.
  • the average particle size of the carbide particles is about 1-30 nm, preferably about 1-15 nm.
  • the average particle size of the carbide particles should not be too large, otherwise it will be detrimental to the reduction of area performance of the steel.
  • the precipitation of a certain amount of carbide particles can significantly reduce the carbon content in the martensite of the steel, thereby improving the performance of the martensite reduction of area.
  • about 0.1-0.5% by weight of carbide particles are precipitated.
  • the total amount of precipitated carbide particles should not be too much, otherwise the coarsening tendency of carbide particles will be significant, which will cause the strength of the steel to decrease and deteriorate the reduction of area performance.
  • N nitrogen
  • the above-mentioned precipitated carbide particles may also contain nitrogen.
  • the carbide particles comprise composite carbide particles of V-Mo, Ti and Nb, which optionally further contain nitrogen.
  • the yield strength of the steel obtained in the present invention is greater than or equal to about 1400 MPa, more preferably greater than or equal to about 1550 MPa; the tensile strength is greater than or equal to about 1800 MPa, preferably greater than or equal to about 1900 MPa; and the reduction of area is greater than or equal to about 38%, preferably greater than or equal to about 45%.
  • the present invention provides a spring member for vehicle suspension made of the above-mentioned steel, including, for example, leaf springs, stabilizer bars, round springs, and the like.
  • the heat-treated steel of the present invention has both high strength, high ductility and high toughness, especially high reduction of area, which is derived from the selection of alloy composition and the selection of heat treatment process conditions.
  • V and Mo elements By introducing V and Mo elements into the high-strength steel used in the present invention, composite carbide particles of V and Mo with controllable average particle size and total amount are formed in the heat-treated steel.
  • the composite carbide particles may further include N.
  • the composite carbide particles may further include Ti and Nb. The formation of composite carbide particles makes the steel have high strength and high reduction of area.
  • the nano-sized carbide particles dispersed in the heat-treated steel have a large surface area, which can be used as a trap point for hydrogen, thereby helping to improve the delayed cracking resistance of the material.
  • V has a higher solid solubility product in austenite than other carbide-forming elements. Therefore, at high temperature, that is, during the austenitizing step, V carbide particles are not easy to precipitate. But when holding at a relatively low temperature, V carbide particles can be precipitated from austenite in large quantities, and the carbide size is fine. Mo is not easy to precipitate in austenite, but it is added together with V to form composite carbides with V. Therefore, the heat treatment method of the present invention introduces a carbide precipitation step.
  • This step ensures that the composite carbide particles of V and Mo or the composite carbide particles of V-Mo and Ti, Nb can be fully analyzed on the austenite grain boundaries and in the grains, and the average of the precipitated carbide particles can be controlled at the same time Particle size and total amount.
  • the dispersion of a large number of nano-scale carbide particles in the heat-treated steel not only increases the strength of the steel but also improves the martensite reduction of area, and is beneficial to the delayed cracking resistance of the material.
  • the low-temperature tempering step further improves the reduction of area performance of the steel.
  • the U-notch impact energy at -40°C is tested in accordance with the "GBT229-2007 Charpy Pendulum Impact Test Method for Metallic Materials".
  • the sample size is 55 ⁇ 10 ⁇ 10mm 3 .
  • the thickness of the decarburized layer is tested according to the microhardness test method of "GB224-2008 Steel Decarburized Layer Depth Measurement Method".
  • the thickness of the decarburized layer is defined as the distance from the surface of the sample to the point where the hardness of the core reaches 50%.
  • the relative proportions of martensite (M), ferrite (F) and bainite (B) are measured by quantitative metallography.
  • the ratio of retained austenite (RA) is tested by XRD.
  • the average particle size and total amount of carbide particles are statistically obtained by shooting 5 fields of view randomly under a transmission electron microscope.
  • the chemical composition of carbides is tested by using the EDS function under a transmission electron microscope.
  • the high-strength steel having the composition shown in Table 1 below was prepared for use in the heat treatment method of the present invention.
  • the high-strength steel is made by heating a billet with the composition shown in Table 1 below to 1200°C for 60 minutes, rolling it at 900°C, and cooling it to room temperature at a cooling rate of 30°C/min. The thickness is 16mm. Hot rolled flat steel.
  • A1-A5 is the high-strength steel of the present invention, and B1-B3 are comparative steels.
  • Table 1 Steel composition (wt%, the balance is Fe and other unavoidable impurities except P and S)
  • the microstructure of the steel obtained by the heat treatment of the invention using the steels A1-A5 of the invention contains more than 93% martensite, less than 4% retained austenite, less than 2% ferrite, and Bainite less than 3%.
  • the average particle size of the composite carbide particles containing V and Mo is 5-15 nm, and the amount of carbide particles is 0.15-0.40%.
  • the yield strength of the steel materials A1-A5 of the present invention can reach 1400-1750 MPa
  • the tensile strength can reach 1850-2150 MPa
  • the reduction of area is 45-60%.
  • the surface decarburization layer can be controlled below 100 ⁇ m. This helps to improve the fatigue performance of the spring member formed therefrom.
  • the C content of the comparative steel B1 used in Comparative Example 1 is lower than the lower limit required by the present invention, and it does not contain Mo and V.
  • the heat treatment does not include the carbide precipitation step and the tempering step of the present invention.
  • the low-carbon design ensures that the steel has a good reduction in area performance.
  • the steel does not contain Mo, Ti, and V elements and the heat treatment does not include the carbide precipitation step, the grain refinement and precipitation strengthening effect are not significant.
  • the yield strength of the comparative steel B1 after heat treatment is only 1214 MPa, and the tensile strength is only 1563 MPa, which does not meet the requirements for preparing a spring member for vehicle suspension.
  • the C content of the comparative steel B2 used in Comparative Example 2 is higher than the upper limit required by the present invention, and it does not contain Mo and V.
  • the heat treatment does not include the carbide precipitation step of the present invention, and the tempering step is a medium-high temperature tempering process whose temperature exceeds the scope of the present invention.
  • the high-carbon design ensures the high strength of the steel, but due to the absence of carbide precipitation steps, high-carbon martensite appears, resulting in a low reduction of area.
  • the tempering process at medium and high temperature in Comparative Example 2 can improve the reduction of area to a certain extent, the martensite is tempered and softened during the tempering process at medium and high temperature, resulting in a significant decrease in strength.
  • the C content and Si content of the comparative steel B3 used in Comparative Example 3 are both higher than the upper limit required by the present invention, and the Mo content is lower than the lower limit required by the present invention.
  • the heat treatment does not include the carbide precipitation step of the present invention, and the tempering step is an intermediate temperature tempering process whose temperature exceeds the scope of the present invention.
  • the high C design ensures the high strength of the steel, and the high Si design stabilizes a large amount of retained austenite in the steel. When plastic deformation occurs, the retained austenite has a TRIP effect, which improves the ductility of the steel.
  • the heat-treated steel material B3 in Comparative Example 3 has high strength, its reduction of area is low.
  • the comparative steel B3 has a high Si content, which causes serious decarburization on the surface of the comparative steel B3 after heat treatment.
  • the thickness of the decarburized layer reaches more than 200 ⁇ m, which will significantly reduce the fatigue performance of the component.
  • the heat-treated steel B3 in Comparative Example 3 is not suitable for the requirements of preparing spring members for vehicle suspensions.
  • Comparative Example 4 uses the steel A1 of the present invention, but the heat treatment does not include the carbide precipitation step of the present invention. Therefore, although carbide particles with a small average particle size are precipitated, the total amount of precipitated carbide particles is relatively small. Correspondingly, the carbon content of martensite is not significantly reduced, resulting in insufficient improvement in the ductility and toughness of the steel. At the same time, the precipitation strengthening effect is not ideal. Therefore, the strength of the heat-treated steel A1 in Comparative Example 4 is low, and the reduction of area is also low, which does not meet the requirements for preparing a spring member for vehicle suspension.

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Abstract

一种高强度钢的热处理方法，其中以重量百分比计，所述高强度钢包含：0.30-0.45％的C，1.0％以下的Si，0.20-2.5％的Mn，0.20-2.0％的Cr，0.15-0.50％的Mo，0.10-0.40％的V，0.2％以下的Ti，0.2％以下的Nb，其余为Fe和其他合金元素及杂质，其中合金成分使得式(1)的Eq(Mn)不小于1.82，Eq(Mn)＝Mn+0.26 Si+3.50 P+1.30 Cr+2.67 Mo (1)。所述方法包括以下步骤：1)奥氏体化步骤；2)碳化物析出步骤；和3)回火步骤。热处理后的钢材同时具有高强度、高延性和高韧性，特别是具有高的断面收缩率，从而特别适用于制备车辆悬架用的弹簧构件。

Description

一种高强度钢的热处理方法和由此获得的产品 技术领域
本发明涉及高强度钢的热处理。更具体地说，本发明涉及使处理后的钢材同时具有高强度、高延性和高韧性的高强度钢热处理方法，这种处理后的钢材特别适用于制备弹簧构件，例如车辆悬架用的弹簧构件。
背景技术
车辆悬架用弹簧构件包括，例如钢板弹簧、稳定杆、圆簧等。钢板弹簧简称板簧，常被安装在车架与车桥之间。在车辆行驶过程中，钢板弹簧承受车轮对车架的载荷冲击，消减车身的剧烈振动，保持车辆行驶的平稳性和对不同路况的适应性。稳定杆是一种扭杆弹簧。车身在转弯时，稳定杆利用杆身的弹力阻止车轮抬起，防止车身发生过大的横向侧倾，尽量使车身保持平衡。在车辆行驶过程中，车辆悬架用弹簧构件反复承受应力的负荷。为保证汽车行驶的平顺性，通常希望将悬架的刚度降低，这进一步提高对车辆悬架用弹簧构件的承受应力负荷的能力要求。因此，制备车辆悬架用弹簧构件所使用的钢材需要具有高强度。
此外，为达到节能减排、降低制造成本，提升车辆安全性能等目标，汽车轻量化的要求越来越高，这进一步要求提高钢材强度。
然而，随着钢材强度的提高，其他性能(特别是延性和韧性)急剧下降，这限制了高强度钢在汽车构件上的应用。
近年来，对于开发可用于制备车辆悬架弹簧构件的高性能钢材，本领域中已经做出了不少尝试。
中国专利申请CN108239726A提供了一种耐氢脆性优良的高强度弹簧用钢材及其制造方法。其中，化学成分以重量百分比计，所述钢材包含：C 0.45-0.60％，Si 1.40-1.80％，Mn 0.30-0.80％，Cr 0.20-0.70％，Mo 0.05-0.15％，V 0.05-0.20％，Nb 0.010-0.030％，N≤0.006％，P≤0.015％，S≤0.015％。相应的热处理工艺为：通过将钢材加热至880-1000℃保温10s-30min，然后以10℃/s以上的平均冷速进行冷却，随后，再将钢材加热至380-460℃保温10s-40min，然后以10℃/s以上的平均冷速冷至 60℃以下。CN108239726A认为，通过添加Nb、V、Mo等强碳化物形成元素，在钢材内析出一定量的平均晶粒直径在10-60nm的(V、Mo)C或(Nb、V、Mo)C颗粒，并且原奥氏体晶粒也细化至10级以上。碳化物颗粒的界面活化能高，可以起到不因外部应力而扩散的非扩散性氢捕集作用。同时，细小的原始奥氏体晶粒以及足够量的纳米级碳化物颗粒保证了弹簧用钢材在抗拉强度在1900MPa以上的同时耐氢脆性优良。
中国专利申请CN106399837A提供了一种热冲压成形用钢材及热成形工艺。其中，化学成分以重量百分比计，钢材包含：C 0.27-0.40％，Si 0-0.80％，Mn 0.20-3.0％，V 0.10-0.4％，Si 0-0.8％，Al 0-0.5％，Cr 0-2％，Ti 0-0.15％，Nb 0-0.15％，B 0-0.004％以及总含量小于2％的Mo、Ni、Cu。通过添加V、Nb等强碳化物形成元素，结合热成形工艺条件，控制纳米级的VC颗粒和/或V与Ti、Nb的复合碳化物颗粒在热成形过程中析出，起到析出强化和晶粒细化的作用，并降低马氏体的碳含量，改善了钢材的韧性。随后，通过低温回火，钢材的性能得到进一步优化。其所得钢材具有1350-1800MPa的屈服强度，1700-2150MPa的抗拉强度，以及7-10％的延伸率。
对于车辆悬架用弹簧构件来说，断面收缩率经常被用作综合表征钢材的延性和韧性的重要指标。为保证钢材具有高的断面收缩率性能，需要钢材同时具有高延性和高韧性。目前公开的钢材所具有的延性和韧性并不足以提供适用于制备车辆悬架用弹簧构件的断面收缩率性能。因此，仍然需求一种高强度钢的热处理方法，其能够使处理后的钢材同时具有高强度、高延性和高韧性，特别是具有高的断面收缩率，从而特别适用于制备车辆悬架用的弹簧构件。
发明内容
本发明提供一种获得具有高的断面收缩率的高强度钢的热处理方法，解决了上述问题。在一个方面，本发明提供了一种高强度钢的热处理方法，其中以重量百分比计，所述高强度钢包含：0.30-0.45％的C，1.0％以下的Si，0.20-2.5％的Mn，0.20-2.0％的Cr，0.15-0.50％的Mo，0.10-0.40％的V，0.2％以下的Ti，0.2％以下的Nb，其余为Fe和其他合金元素及杂质，其中合金成分使得式(1)的Eq(Mn)不小于1.82，
Eq(Mn)＝Mn+0.26Si+3.50P+1.30Cr+2.67Mo       (1)
所述方法包括以下步骤：
1)奥氏体化步骤：将所述高强度钢加热至约奥氏体化临界温度(Ac3)以上20℃(以下简称Ac3+20℃)至约950℃，保温约1-300min；
2)碳化物析出步骤：在奥氏体化步骤后，将所述高强度钢初始冷却至约铁素体析出开始温度(Ar3)以下10℃(以下简称Ar3-10℃)至约870℃，保温约5-300min，随后进一步冷却至约100℃以下，其中所述进一步冷却的平均冷速不小于约1℃/s；和
3)回火步骤：在碳化物析出步骤后，将所述高强度钢再次加热至约120-280℃，保温约5-360min。
在另一个方面，本发明提供经过上述热处理方法获得的钢材，其中以面积计，所述的钢材的显微组织包括：大于等于约90％的马氏体，小于等于约3％的铁素体，小于等于约5％的残余奥氏体，小于等于约10％的贝氏体。
其中所述钢材含有约0.1-0.5重量％的碳化物颗粒，其中所述碳化物颗粒包含V与Mo的复合碳化物颗粒，并且所述碳化物颗粒的平均粒径为约1-30nm，并且
其中所述钢材的屈服强度大于等于约1400MPa，抗拉强度大于等于约1800MPa，和断面收缩率大于等于约38％。
在又一个方面，本发明提供由上述钢材制备的车辆悬架用弹簧构件。
附图说明
附图用以辅助对本发明的进一步理解，并且构成说明书的一部分，与下面的具体实施方式一起用于说明本发明，但并不构成对本发明的限制。在附图中：
图1为本发明的高强度钢的热处理方法的一个实施方案所用的温度-时间示意图；
图2为经过本发明的热处理方法获得的钢材的一个实施方案的金相照片；
图3为经过本发明的热处理方法获得的钢材的一个实施方案的透射电镜照片；
图4为经过本发明的热处理方法获得的钢材的一个实施方案的碳化物化学成分结果。
具体实施方式
除了在实施例中之外，本说明书中的参数(例如，量或条件)的所有数值应理解为在所有情况下都由术语“约”修饰，无论“约”是否实际出现在数值之前。
除非特别指出，否则说明书中所提到的百分比为重量百分比。
根据本发明的一个方面，涉及一种高强度钢的热处理方法，其中以重量百分比计，所述高强度钢包含：0.30-0.45％的C，1.0％以下的Si，0.20-2.5％的Mn，0.20-2.0％的Cr，0.15-0.50％的Mo，0.10-0.40％的V，0.20％以下的Ti，0.2％以下的Nb，其余为Fe和其他合金元素及杂质，其中合金成分使得式(1)的Eq(Mn)不小于1.82，
Eq(Mn)＝Mn+0.26Si+3.50P+1.30Cr+2.67Mo     (1)
所述方法包括以下步骤：
1)奥氏体化步骤：将所述高强度钢加热至约Ac3+20℃至约950℃，保温约1-300min；优选将所述高强度钢加热至约Ac3+30℃至910℃，保温约1-30min；
2)碳化物析出步骤：在奥氏体化步骤后，将所述高强度钢初始冷却至约Ar3-10℃至约870℃，保温约5-300min，随后进一步冷却至约100℃以下，其中所述进一步冷却的平均冷速不小于约1℃/s；优选将所述高强度钢初始冷却至约Ar3+10℃至850℃保温约5-30min；和
3)回火步骤：在碳化物析出步骤后，将所述高强度钢再次加热至约120-280℃，保温约5-360min；优选所述高强度钢再次加热至约160-230℃，保温约10-60min。
以重量百分比计，优选所述高强度钢包含：0.32-0.42％的C，0.8％以下的Si，0.2-1.5％的Cr，0.2-0.4％的Mo，0.12-0.3％的V，其余为Fe和其他合金元素及杂质，其中合金成分使得式(1)的Eq(Mn)不小于1.82。
本发明所用的高强度钢的成分详细描述如下。
C：约0.30-0.45％
C是钢中最有效的固溶强化元素，为保证钢材的抗拉强度在 1800MPa以上，需要C含量大于等于约0.30％。但是，如果C含量超过0.45％，则形成的高碳马氏体，其延性和韧性均较差，且抗氢脆性能显著下降。因此，本发明所用的高强度钢的C含量在约0.30-0.45％之间；优选为约0.32-0.42％之间。
Si：约1.0％以下
Si是钢材熔炼时的脱氧剂，且固溶在铁素体基体中具有强化母材强度的作用。然而，过量的Si不仅对钢材韧性有害，且在热处理过程中会形成较严重的表面氧化和脱碳。脱碳层厚度为车辆悬架用弹簧构件疲劳性能的关键控制参数之一。因此，本发明所用的高强度钢的Si含量为1.0％以下，优选为0.8％以下。
Mn：约0.20-2.5％
Mn是提高钢材淬透性确保强度的元素。当Mn含量低于约0.20％时，钢材的淬透性不足，难以获得高强度。但是，当Mn含量过高时，钢材的延性和韧性将显著下降，因此，本发明所用的高强度钢的Mn含量上限为约2.5％。
Cr：约0.20-2.0％
Cr是提高钢材淬透性的元素，同时对钢材的抗氧化具有显著效果。在Cr含量不足约0.20％的情况下，上述效果不明显。相反，当Cr含量超过2.0％时，易析出粗大含Cr碳化物，对韧性不利。因此，本发明所用的高强度钢的Cr含量范围为约0.20-2.0％，优选为0.2-1.5％。
Mo：约0.15-0.50％
Mo是强碳化物形成元素，与碳原子的亲和力较大，能够阻止碳原子的扩散，降低碳元素的扩散系数，进而有效抑制钢材的表面脱碳问题。脱碳层厚度为车辆悬架用弹簧构件疲劳性能关键控制参数之一。同时，Mo的添加也同时提高了钢材的淬透性，且恰好利用了Mo与V在热处理过程中析出复合纳米碳化物。复合碳化物的析出更有利于碳化物的弥散分布以及获得更细小的碳化物。这既保证了钢材的超高强度也使得材料具有良好的断面收缩率性能。因此，本发明的钢材中添加的Mo含量不低于约0.15％。但如果Mo含量高于约0.50％，则会导致生产成本显著上升。因此，Mo含量为约0.15-0.50％，优选Mo的含量为大于约0.20至约0.40％。当Mo的含量为大于约0.20至约0.40％时，可以有效抑制或缓解弹簧构件的表面脱碳问题，从而为弹簧构件 提供良好的耐疲劳性能，同时该范围的Mo的添加也能确保获得Mo与V的复合碳化物弥散分布且尺寸细小，从而为弹簧构件提供良好的断面收缩率。
V：约0.10-0.40％
V形成复合碳化物，从而起到析出强化和原始奥氏体晶粒细化的作用。如果V含量低于约0.10％，则无法形成充足的碳化物，上述效果不显著。如果V含量高于约0.40％，则会导致生产成本上升和碳化物粗大导致断面收缩率下降。因此，优选V含量为约0.10-0.40％，优选为0.12-0.30％。
Ti：小于等于约0.20％，Nb：小于等于约0.20％
Ti和Nb在钢中均形成碳氮化物，对强度的提高和晶粒的细化均有影响。因为Ti和Nb是最强的碳化物形成元素，所以当含量超过约0.20％时，则在高温下会大量析出碳氮化物。这导致其尺寸粗大，导致断面收缩率下降。若通过热处理控制Ti或Nb在高温下析出碳氮化物，以尽量少在高温下析出碳氮化物，则有利于其与V和Mo的复合析出，可以进一步细化碳氮化物的尺寸。但是，这样的工艺控制较为复杂。因此，Ti的含量为小于等于约0.20％，优选为小于等于0.05％；Nb含量为小于等于约0.20％，优选为小于等于约0.05％。
如上所述，Ti和Nb都具有在钢中形成碳氮化物而提高钢材综合性能的效果。发明人还发现，通过将两者复合添加会起到协同效果。因此，Ti+Nb的合计总量为约0.20％以下。如果Ti+Nb的合计总量超过约0.20％，则碳氮化物在高温即会大量析出，导致其尺寸粗大，导致断面收缩率下降。因此，Ti+Nb的合计总量小于等于约0.20％，优选为小于等于约0.08％。
P、S：分别小于等于约0.025％
P和S会在晶界偏析，从而导致钢材断面收缩率下降。因此，希望这些元素尽可能少，例如P和S的含量均小于等于约0.025％。
Eq(Mn)：不小于约1.82
Eq(Mn)表征了钢材的淬透性。Eq(Mn)与马氏体的临界冷速Rc的关系满足下式(2)。
lgRc＝3.15-1.73Eq(Mn)      (2)
为保证钢材具有良好的淬透性，即Rc小于等于约1℃/s，则需保 证Eq(Mn)不小于约1.82。
本发明的余量成分是铁(Fe)。但是，在常规制造过程中来自原料或周围环境的杂质会不可避免地被混入，因此，不能排除这些杂质被混入。这些杂质是所属领域的普通技术人员都知道的。
如图1所示，本发明的热处理方法包括对上述高强度钢实施以下步骤：1)奥氏体化步骤；2)碳化物析出步骤；和3)回火步骤。
具体地说，在本发明的热处理中，高强度钢经历奥氏体化步骤。奥氏体化步骤通过将所述高强度钢加热至约Ac3+20℃至约950℃，保温约1-300min来进行。当加热温度低于Ac3+20℃时或保温时间不足1min时，则可能会有未固溶的铁素体和珠光体残留，导致合金元素未均匀化。当加热温度高于950℃或保温时间超过60min时，则可能导致钢材表面氧化和脱碳严重，并且奥氏体的晶粒粗大。因此，将奥氏体化步骤的工艺条件设定为在Ac3+20℃-950℃保温1-60min，优选在Ac3+30℃-910℃保温1-30min。通常奥氏体化步骤通过在炉温为Ac3+40℃至约970℃的加热炉中保温5-310min实现，优选在炉温为Ac3+50℃至约930℃的加热炉中保温5-40min实现。另外，上述奥氏体化步骤也可应用感应加热或感应加热与加热炉加热相结合的方式完成。
在奥氏体化步骤后，所述高强度钢经历碳化物析出步骤。碳化物析出步骤如下进行：将所述高强度钢初始冷却至约Ar3-10℃至于约870℃，保温约5-300min；优选将所述高强度钢初始冷却至约Ar3+10℃至850℃保温5-30min，；随后进一步冷却至约100℃以下，其中所述进一步冷却的平均冷速大于等于约1℃/s，优选大于等于约5℃/s。所述进一步冷却的平均冷速可以小于等于约100℃/s，优选小于等于约50℃/s，更优选小于等于约20℃/s。所述初始冷却可在炉温为约Ar3-20℃至约870℃的加热炉中保温约5-300min实现，优选在炉温为约Ar3℃至于约850℃的加热炉中保温约5-30min实现。所述进一步冷却可以通过包括油淬、盐水淬火等的方式来进行。当所述的碳化物析出步骤的初始冷却的保温温度低于Ar3-10℃时，可能有较多的铁素体生成，这将对钢材的强度以及疲劳性能有害。当所述保温温度高于870℃时，碳化物析出困难，且碳化物会粗化。当保温时间小于5min时，碳化物析出不充分，而当保温时间超过300min，可能导致奥氏体晶粒长大和碳化物粗 化，从而对钢材的断面收缩率性能不利。而且，较长时间的保温也导致了钢材表面氧化和脱碳严重。因此，将碳化物析出步骤的工艺条件设定为在约Ar3-10℃-870℃保温约5-300min，随后进一步冷却至约100℃以下。更优选在约Ar3+10℃-850℃保温约5-60min，随后进一步冷却至约100℃以下。在一个方案中，进一步冷却至约80℃或更低，例如室温。为保证冷却过程中尽可能少地生成铁素体而尽可能多地生成马氏体，冷却过程的平均冷速设定为大于等于约1℃/s。上述碳化物析出步骤可与奥氏体化步骤在同一加热炉内完成，亦可以在同一加热炉的不同分段完成，亦可以在不同的加热炉中完成，亦可以通过任意其它加热方式完成。优选的是，本发明处理的高强度钢的初始组织中存在大量纳米碳化物颗粒。因此，在奥氏体化步骤后，所述高强度钢的组织中仍有大量纳米碳化物颗粒未溶解，这有利于在碳化物析出步骤中对碳化物颗粒的析出量和尺寸进行控制。
在碳化物析出步骤后，所述高强度钢经历回火步骤。回火步骤通过将所述高强度钢再次加热至约120-280℃，保温约5-360min来进行。当回火温度低于约120℃或保温时间小于约5min时，马氏体的回火效果不充分，马氏体因相变导致的内应力得不到充分释放，其断面收缩率性能无法得到进一步改善。当回火温度高于约280℃或保温时间超过约360min时，将导致Fe-C碳化物的大量析出，将导致钢材的强度下降明显且碳化物粗化趋势显著。优选为在温度约120-250℃的炉中加热约10-120min。通常，汽车构件加工制造过程中包括在涂装后进行烘烤的步骤。该烘烤通过在约150-230℃保温约10-60分钟来进行。该烘烤步骤可以起到上述回火步骤的作用，从而无需额外的回火步骤。
经过上述热处理方法，获得了一种钢材，其中以面积计，所述钢材的显微组织包括：大于等于约90％的马氏体，小于等于约3％的铁素体，小于等于约5％的残余奥氏体，小于等于约10％的贝氏体；优选包括大于等于约97％的马氏体，并且残余奥氏体、铁素体与贝氏体的总和小于等于约2.5％；
其中所述钢材含有约0.1-0.5重量％的碳化物颗粒，其中所述碳化物颗粒包含V与Mo的复合碳化物颗粒，并且所述碳化物颗粒的平均粒径为约1-30nm，并且
其中所述钢材的屈服强度大于等于约1400MPa，抗拉强度大于等 于约1800MPa，和断面收缩率大于等于约38％；优选所述钢材的屈服强度大于等于约1550MPa，抗拉强度大于等于约1900MPa，和断面收缩率大于等于约45％。
本发明热处理方法所获得的钢材详细描述如下。
在本发明获得的钢材中，显微组织包括大于等于约90％的马氏体。马氏体是为得到高强度所需要的显微组织。当马氏体的面积百分数小于约90％，意味着对强度上升贡献较小的铁素体和残余奥氏体过多，从而难以实现高的抗拉强度。本发明获得的钢材中，显微组织包括大于等于约90％的马氏体，从而保证了钢材的强度。所述马氏体的面积百分数优选为大于等于约97％，并可以为约99％以上。
在本发明获得的钢材中，显微组织包括小于等于约10％的贝氏体。贝氏体比马氏体的硬度低。相应地，贝氏体在钢材中的存在会降低钢材的强度。因此，贝氏体含量不宜超过10％。优选贝氏体含量小于等于约3％，并可以为约0％。
在本发明获得的钢材中，显微组织包括小于等于约3％的铁素体。铁素体为软相。当其与马氏体共同存在于钢材中时，两者形成较大的硬度差，这会显著降低钢材的强度。因此，应尽可能避免铁素体的产生，优选为铁素体含量小于等于约1％，并可以为约0％。。
在本发明获得的钢材中，显微组织包括小于等于约5％的残余奥氏体。残余奥氏体可以提高钢材的延性，且有助于提高钢材的耐氢脆性。因此，本发明的钢材中可以包含一定量的残余奥氏体。但残余奥氏体会降低钢材强度，因此不宜过多，过量的残余奥氏体在钢材发生塑性变形时会形成高碳马氏体，从而对钢材的韧性有害。优选残余奥氏体小于等于3％，更优选小于等于1％。
在本发明获得的钢材含有约0.1-0.5重量％的碳化物颗粒，其中所述碳化物颗粒包含V与Mo的复合碳化物颗粒，并且所述颗粒的平均粒径为约1-30nm。在碳化物析出步骤中，碳化物颗粒会在奥氏体晶界上析出，其钉扎奥氏体晶粒，从而抑制了奥氏体晶粒的长大。同时，碳化物颗粒也在奥氏体晶粒内析出。在奥氏体晶粒内析出的是更均匀更细小的第二相粒子，这可以起到析出强化的作用，使得钢材强度提高。所述碳化物颗粒的平均粒径为约1-30nm，优选约1-15nm。所述碳化物颗粒的平均粒径不宜过大，否则对钢材的断面收缩率性能不利。 同时，一定量的碳化物颗粒析出可显著降低钢材马氏体中的碳含量，由此改善马氏体的断面收缩率性能。在本发明获得的钢材中，有约0.1-0.5重量％的碳化物颗粒析出。析出的碳化物颗粒的总量不宜过多，否则碳化物颗粒的粗化趋势显著，反而导致钢材的强度下降且恶化断面收缩率性能。
钢材中不可避免会含有一定量的氮(N)，通常为约0.002-0.008％。因此上述析出碳化物颗粒极可能也含有氮。
在另一个方案中，所述碳化物颗粒包含V-Mo与Ti、Nb的复合碳化物颗粒，其任选进一步含有氮。
在本发明获得的钢材的屈服强度大于等于约1400MPa，更优选大于等于约1550MPa；抗拉强度大于等于约1800MPa，优选约大于等于约1900MPa；和断面收缩率大于等于约38％，优选大于等于约45％。
在又一个方面，本发明提供由上述钢材制备的车辆悬架用弹簧构件，包括例如钢板弹簧、稳定杆、圆簧等。
无需束缚于任何理论，据信本发明热处理后的钢材同时具有高强度、高延性和高韧性，特别是高断面收缩率，源自于对合金成分的选择，配合对热处理工艺条件的选择。通过将V、Mo元素引入本发明所用的高强度钢，在热处理后的钢材中形成平均粒径和总量可控的V与Mo的复合碳化物颗粒。所述复合碳化物颗粒可以进一步包含N。所述复合碳化物颗粒还可以进一步包含Ti和Nb。复合碳化物颗粒的形成使得钢材具有高强度和高断面收缩率。此外，弥散在热处理后的钢材中的纳米级碳化物颗粒具有极大的表面积，可以作为氢的捕获点，从而有利于改善材料的抗延迟开裂性能。V在奥氏体中较其他碳化物形成元素具有更高的固溶度积。因此，在高温时，即奥氏体化步骤中，V的碳化物颗粒不易析出。但当在相对低温度保温时，V的碳化物颗粒可以从奥氏体中大量析出，且碳化物尺寸细小。Mo在奥氏体中不易析出，但与V同时添加，可与V形成复合碳化物。因此，本发明的热处理方法引入了碳化物析出步骤。该步骤保证了V与Mo的复合碳化物颗粒或V-Mo与Ti、Nb的复合碳化物颗粒在奥氏体晶界上和晶粒内可以充分析出，同时可控制析出的碳化物颗粒的平均粒径和总量。在热处理后的钢材中弥散大量纳米级的碳化物颗粒既提高了该钢材的强度又改善马氏体的断面收缩率性能，且对材料的抗延迟开裂性能有益。 同时，低温回火步骤使得所述钢材的断面收缩率性能得到进一步的提升。
性能测试
拉伸性能
拉伸性能是按照《GBT 228.1-2010金属材料拉伸试验第1部分：室温试验方法》进行测试，采用R6圆棒状试样。
-40℃的U型缺口冲击功
-40℃的U型缺口冲击功是按照《GBT229-2007金属材料夏比摆锤冲击试验方法》进行测试，试样尺寸为55×10×10mm 3。
脱碳层厚度
脱碳层厚度是按照《GB224-2008钢的脱碳层深度测定法》的显微硬度测试方法进行测试。脱碳层厚度规定为从试样表面至达到50％心部硬度的点的距离。
显微组织特征
马氏体(M)、铁素体(F)和贝氏体(B)的相比例采用定量金相法测定。残余奥氏体(RA)的相比例通过XRD进行测试。碳化物颗粒的平均粒径和总量是通过在透射电镜下随机拍摄5个视场后进行统计获得。碳化物的化学成分是通过在透射电镜下采用EDS功能进行测试。
实施例
通过以下实施例使明本发明的特点和优点显而易见。实施例旨在描述而非以任何方式限制本发明。
准备具有下表1所示组成的高强度钢，以用于进行本发明的热处理方法。所述高强度钢是将具有下表1所示组分的钢坯加热至1200℃保温60min并在900℃下进行轧制后以30℃/min的冷却速度冷却至室温而制造的厚度为16mm的热轧扁钢。A1-A5本发明的高强度钢，B1-B3为对比钢材。
表1钢材组成(wt％，余量为Fe和除了P、S以外的其他不可避免的杂质)
Figure PCTCN2019111796-appb-000001
将上述钢材A1-A5和B1-B3分别按照图1示意的工艺或类似工艺进行热处理，其中热处理的各步骤的工艺条件如下表2显示。
表2热处理的工艺条件
Figure PCTCN2019111796-appb-000002
如上文所述对热处理后的钢材的拉伸性能、-40℃的U型缺口冲击功、显微组织特征进行测定，测试结果见表3。图2显示了实施例1的热处理后的钢材A1的金相照片，其中热处理后的钢材A1的显微组织主要为马氏体。图3显示了实施例1的热处理后的钢材A1中碳化物颗粒的平均粒径为8.2nm，含量约为0.27％。由图4可知，碳化物颗粒包含Mo、V和Nb。
表3热处理后的钢材的性能与显微组织特征的测量结果
Figure PCTCN2019111796-appb-000003
如上表3显示，使用本发明钢材A1-A5进行的本发明热处理获得的钢材的显微组织包含93％以上的马氏体，4％以下的残余奥氏体，2％以下的铁素体以及3％以下的贝氏体。同时，其含有V与Mo的复合碳化物颗粒的平均粒径为5-15nm，碳化物颗粒的量为0.15-0.40％。相应地，通过本发明热处理后的本发明钢材A1-A5的屈服强度可达1400-1750MPa，抗拉强度可达1850-2150MPa，断面收缩率为45-60％。 这说明，通过对合金成分的选择，配合对热处理工艺条件的选择，热处理后的本发明钢材中弥散析出大量的纳米级碳化物颗粒，这降低了马氏体内的碳含量，由此保证马氏体具有高断面收缩率。同时，碳化物颗粒带来的晶粒细化和析出强化效果，对钢材的高强度做出进一步贡献。因此，使用本发明热处理的本发明钢材A1-A5均获得优异的综合性能，完全符合车辆悬架用弹簧构件制备要求。此外，本发明钢材中添加有Mo同时具有较低C含量。本发明钢材经过本发明热处理后，其表面脱碳情况得到有效抑制。表面脱碳层厚度可控制在100μm以下。这更有助于改善由其形成的弹簧构件的疲劳性能。
对比例1使用的对比钢材B1的C含量低于本发明要求的下限，同时其并不含有Mo和V。另外，热处理也未包括本发明的碳化物析出步骤和回火步骤。低碳设计保证了钢材具有良好的断面收缩率性能。但是，由于钢材中不含Mo、Ti和V元素并且热处理中不包括碳化物析出步骤，晶粒细化以及析出强化作用不显著。相应地，热处理后的对比钢材B1的屈服强度仅为1214MPa，抗拉强度仅为1563MPa，并不符合制备车辆悬架用弹簧构件的要求。
对比例2使用的对比钢材B2的C含量高于本发明要求的上限，同时其并不含有Mo和V。另外，热处理并不包括本发明的碳化物析出步骤，并且回火步骤为温度超过本发明范围的中高温回火工艺。高碳设计保证了钢材的高强度，但由于无碳化物析出步骤，出现高碳马氏体，导致断面收缩率偏低。虽然对比例2采用中高温的回火工艺可以在一定程度上改善钢材的断面收缩率，但是在中高温回火过程中，马氏体发生回火软化，导致强度显著降低。同时，中高温回火过程中还有渗碳体析出，从而对钢材韧性有害。因此，对比例2中热处理后的钢材B2的强度和断面收缩率均偏低，并不符合制备车辆悬架用弹簧构件的要求。
对比例3使用的对比钢材B3的C含量和Si含量均高于本发明要求的上限，Mo含量低于本发明要求的下限。另外，热处理并不包括本发明的碳化物析出步骤，并且回火步骤为温度超过本发明范围的中温回火工艺。高C设计保证了钢材的高强度，高Si设计使得钢材中稳定了较多量的残余奥氏体，当发生塑性变形时，残余奥氏体发生了TRIP效应，改善了钢材的延性。但是，在中温回火工艺中，从马氏体中析 出的碳化物颗粒的粗化趋势显著且含量过多，这反而对材料韧性不利。此外，钢材中稳定的较多的残余奥氏体在发生塑性变形时，发生了TRIP效应生成高碳马氏体，进一步损害钢材的韧性。因此，尽管对比例3中热处理后的钢材B3具有高强度，但其断面收缩率偏低。同时，对比钢材B3具有高Si含量，这造成热处理后的对比钢材B3表面脱碳严重。脱碳层厚度达到200μm以上，这将显著降低构件的疲劳表现。综上，对比例3中热处理后的钢材B3并不适合制备车辆悬架用弹簧构件的要求。
对比例4采用了本发明钢材A1，但热处理并不包括本发明的碳化物析出步骤。因此，尽管析出平均粒径小的碳化物颗粒，但析出的碳化物颗粒的总量偏少。相应地，马氏体的含碳量降低不够显著，导致钢材的延性和韧性改善不足。同时，析出强化效果也不理想。因此，对比例4中热处理后的钢材A1的强度偏低，断面收缩率也偏低，并不符合制备车辆悬架用弹簧构件的要求。
以上描述了本发明的优选实施例，但是本领域的技术人员应该明白的是，在不脱离本发明构思的前提下进行的任何可能的变化或替换，均属于本发明的保护范围。

Claims (11)

  1. 一种高强度钢的热处理方法，其中以重量百分比计，所述高强度钢包含：0.30-0.45％的C，1.0％以下的Si，0.20-2.5％的Mn，0.20-2.0％的Cr，0.15-0.50％的Mo，0.10-0.40％的V，0.2％以下的Ti，0.2％以下的Nb，其余为Fe和其他合金元素及杂质，其中合金成分使得式(1)的Eq(Mn)不小于1.82，
    Eq(Mn)＝Mn+0.26Si+3.50P+1.30Cr+2.67Mo  (1)
    所述方法包括以下步骤：
    1)奥氏体化步骤：将所述高强度钢加热至Ac3+20℃至950℃，保温1-300min；
    2)碳化物析出步骤：在奥氏体化步骤后，将所述高强度钢初始冷却至Ar3-10℃至870℃，保温5-300min，随后进一步冷却至100℃以下，其中所述进一步冷却的平均冷速不小于1℃/s；和
    3)回火步骤：在碳化物析出步骤后，将所述高强度钢再次加热至120-280℃，保温5-360min。
  2. 根据权利要求1所述的方法，其中所述高强度钢包含0.20重量％以下的Ti和Nb。
  3. 根据权利要求1所述的方法，其中奥氏体化步骤如下进行：将所述高强度钢加热至Ac3+30℃-910℃，保温1-30min。
  4. 根据权利要求1所述的方法，其中碳化物析出步骤如下进行：将所述高强度钢初始冷却至Ar3+10℃-850℃保温5-60min，随后进一步冷却至100℃以下。
  5. 根据权利要求1所述的方法，其还包括，在奥氏体化步骤前，将高强度钢成形成预制件。
  6. 通过根据权利要求1所述的方法获得的钢材，其中以面积计，所述的钢材的显微组织包括：大于等于90％的马氏体，小于等于3％的铁素体，小于等于5％的残余奥氏体，小于等于10％的贝氏体，
    其中所述钢材含有0.1-0.5重量％的碳化物颗粒，其中所述碳化物颗粒包含V与Mo的复合碳化物颗粒，并且所述碳化物颗粒的平均粒径为1-30nm，并且
    其中所述钢材的屈服强度大于等于1400MPa，抗拉强度大于等于 1800MPa，和断面收缩率大于等于38％。
  7. 根据权利要求6所述的钢材，其中所述碳化物颗粒进一步含有N。
  8. 根据权利要求6所述的钢材，其中所述碳化物颗粒的平均粒径为约1-15nm。
  9. 根据权利要求6所述的钢材，其中所述钢材的屈服强度大于等于1550MPa；抗拉强度为大于等于1900MPa；和断面收缩率为大于等于45％。
  10. 一种车辆悬架用弹簧构件，其由权利要求6所述的钢材制备。
  11. 根据权利要求10所述的车辆悬架用弹簧构件，其为钢板弹簧、稳定杆或圆簧。
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CN113322365A (zh) * 2021-05-19 2021-08-31 北京理工大学 一种同时提高低碳低合金钢强度和塑性的方法
CN113322365B (zh) * 2021-05-19 2022-05-20 北京理工大学 一种同时提高低碳低合金钢强度和塑性的方法
CN113699448A (zh) * 2021-08-27 2021-11-26 中冶陕压重工设备有限公司 一种低合金结构钢SY41CrMnMoNbVTi及其制备方法
CN114395729A (zh) * 2021-12-13 2022-04-26 唐山中厚板材有限公司 Nm450级无需淬火热处理的耐磨钢板及其生产方法
CN114395729B (zh) * 2021-12-13 2023-09-01 唐山中厚板材有限公司 Nm450级无需淬火热处理的耐磨钢板及其生产方法
CN114457212A (zh) * 2021-12-28 2022-05-10 河钢股份有限公司 一种高温轴承钢碳化物细质弥散处理工艺

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