WO2020078472A1 - 一种800MPa级热冲压桥壳钢及其制造方法 - Google Patents

一种800MPa级热冲压桥壳钢及其制造方法 Download PDF

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WO2020078472A1
WO2020078472A1 PCT/CN2019/111982 CN2019111982W WO2020078472A1 WO 2020078472 A1 WO2020078472 A1 WO 2020078472A1 CN 2019111982 W CN2019111982 W CN 2019111982W WO 2020078472 A1 WO2020078472 A1 WO 2020078472A1
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Prior art keywords
axle housing
mpa
hot stamping
housing steel
steel
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English (en)
French (fr)
Inventor
刘刚
张华伟
王巍
陆敏
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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 DE112019005199.7T priority Critical patent/DE112019005199T5/de
Priority to KR1020217008010A priority patent/KR102495857B1/ko
Publication of WO2020078472A1 publication Critical patent/WO2020078472A1/zh
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D22/00—Shaping without cutting, by stamping, spinning, or deep-drawing
    • B21D22/02—Stamping using rigid devices or tools
    • B21D22/022—Stamping using rigid devices or tools by heating the blank or stamping associated with heat treatment
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D53/00—Making other particular articles
    • B21D53/88—Making other particular articles other parts for vehicles, e.g. cowlings, mudguards
    • B21D53/90—Making other particular articles other parts for vehicles, e.g. cowlings, mudguards axle-housings
    • 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
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • 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
    • C21D6/02—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
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/021—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving particular fabrication steps or treatments of ingots or slabs
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0447—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing characterised by the heat treatment
    • C21D8/0463—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for drawing, e.g. for deep-drawing characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • 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
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the invention relates to a steel type and a manufacturing method thereof, in particular to an 800MPa hot stamping axle housing steel and a manufacturing method thereof.
  • the axle housing As a key load-bearing component of automobiles, the axle housing has high requirements for safety and needs to meet strict component fatigue performance. That is to say, for the steel used for axle housings, stable performance of the steel plate, low temperature impact resistance and good weldability are required. Segregation and inclusions are strictly controlled.
  • China's hot stamping axle housing steels are mainly 16Mn, Q345C, Q420C, Q460C and other common C-Mn structural steels.
  • the strength of such C-Mn steel after hot stamping further decreases, for example, the yield strength of Q460C after hot stamping drops to about 400MPa.
  • axle housing steel which can reach the strength level of 800 MPa, and at the same time, the plasticity and fatigue performance of the steel plate are also good, which is more suitable for the manufacture of axle housings.
  • One of the objects of the present invention is to provide a 800 MPa class hot stamping axle housing steel, which can reach 800 MPa strength level, and at the same time, the 800 MPa class hot stamping axle housing steel also has good plasticity and fatigue performance, and is very suitable for manufacturing axles shell.
  • the present invention proposes an 800MPa hot stamping axle housing steel, the chemical element mass percentage is:
  • C In the 800 MPa hot stamping axle housing steel described in the present invention, carbon plays a role of solid solution strengthening. Adding C can increase the strength of lower bainite. In addition, C can also be used in the hot stamping air cooling process of the axle housing. Fe reacts to form dispersed Fe 3 C, which increases the strength of the steel sheet after hot stamping. However, if the mass percentage of C is too high, it is not conducive to the weldability of the steel plate. Therefore, in the 800 MPa grade hot stamping axle housing steel described in the present invention, the mass percentage of C is controlled to be 0.15 to 0.21%.
  • Si In the 800 MPa grade hot stamping axle housing steel described in the present invention, Si can suppress the precipitation of cementite at high temperature, which is beneficial to the formation of lower bainite, and at the same time, Si can also form fine infiltration in the steel plate Carbon body particles to improve the strength of lower bainite. However, too much Si content is not conducive to the weldability of the steel plate. Therefore, in the technical solution of the present invention, the mass percentage of Si is controlled to be 0.30 to 0.80%.
  • Mn For the technical solution described in the present invention, the addition of a certain content of Mn element is beneficial to promote the formation of lower bainite, and Mn can play a certain solid solution strengthening effect on the lower bainite structure.
  • Mn element is beneficial to the formation of finer ferrite or bainite during hot stamping of the steel sheet, which is beneficial to increase the strength of the steel sheet after hot stamping.
  • an excessively high content of Mn is not conducive to the weldability of the steel sheet. Therefore, the mass percentage of Mn in the 800 MPa grade hot stamping axle housing steel described in the present invention is 1.75-2.10%.
  • Nb For the 800 MPa hot stamping axle housing steel described in the present invention, during the controlled rolling and cooling stage, a small amount of Nb in the steel plate reacts with C to form fine NbC particles, which is conducive to refine the steel plate structure and at the same time It can also improve the strength, plasticity and toughness of the steel plate. In addition, NbC precipitates in the lower bainite structure, which can produce a strong precipitation strengthening effect. In addition, NbC can also refine the austenite grains during the hot stamping heating stage and improve the strength of the axle housing steel sheet after hot stamping. However, when the Nb element content is too high, the NbC particles are large, and the effect of suppressing the growth of austenite grains is weakened. Therefore, the mass percentage of Nb in the 800 MPa grade hot stamping axle housing steel described in the present invention is controlled at 0.015 to 0.040%.
  • Ti element will react with C during the phase transformation of austenite to ferrite after the steel sheet is rolled to form TiC particles with a diameter of several nanometers to several tens of nanometers, thereby producing precipitation strengthening Effect, especially in the range of 680 ⁇ 730 °C can produce finer interphase precipitation.
  • the fine TiC particles in the hot stamping heating stage can suppress the growth of austenite grains, and then refine the structure after hot stamping, and improve the strength of the bridge shell steel sheet after hot stamping.
  • the mass percentage of Ti in the 800 MPa hot stamping axle housing steel described in the present invention is controlled at 0.020 to 0.060%.
  • B In the 800 MPa grade hot stamped axle housing steel described in the present invention, a small amount of B can promote the formation of bainite. However, when the B element content is too high, B brittleness problems are likely to occur, which deteriorates the impact toughness of the steel plate. In addition, in the hot stamping stage of the axle housing, trace elements of B are beneficial to promote the formation of finer bainite and improve the strength of the steel plate. Therefore, the mass percentage of B in the 800 MPa grade hot stamping axle housing steel described in the present invention is controlled at 0.0015 to 0.0030%.
  • Al is an important deoxidizer, usually 0.02% or more of Al is added.
  • the oxide chain inclusions of Al need to be controlled, so the Al content is controlled in a lower range. Therefore, the mass percentage of Al in the 800 MPa hot stamping axle housing steel described in the present invention is controlled to 0.005 to 0.015%.
  • Ca For the 800 MPa hot stamping axle housing steel described in this invention, trace elements of Ca can act as a scavenger in the steel smelting process, improving the toughness and fatigue properties of the steel; meanwhile, Ca treatment can improve MnS The shape of the inclusions prevents the formation of elongated MnS inclusions. However, Ca content exceeding 0.001% tends to form larger size Ca compounds, but deteriorates toughness and fatigue performance. Therefore, the mass percentage of Ca in the 800 MPa hot stamping axle housing steel described in the present invention is controlled at 0.0004 to 0.001%.
  • N For the 800 MPa hot stamping axle housing steel described in the present invention, the N element needs to be controlled in a narrow range in this case. Trace elements of N can react with Ti to form TiN particles, which can effectively suppress the growth of austenite grains during welding and hot stamping, refine the welding heat affected zone and the structure after hot stamping, and improve the heat affected zone and heat The strength, low temperature toughness and fatigue performance of stamped steel plates. However, when the N content is too high, the formed TiN particles are too large, which will deteriorate the low temperature toughness and fatigue performance of the steel plate. Therefore, the mass percentage of N in the 800 MPa grade hot stamping axle housing steel described in the present invention is controlled at 0.001 to 0.004%.
  • P, S, and O are inevitable harmful impurity elements in steel materials, which are not conducive to steel performance.
  • P as an impurity element is prone to cold and brittle problems; S is easy to react with Mn to produce MnS inclusions, which is not conducive to fatigue performance in steel; O It is easy to react with Al to produce Al 3 O chain inclusions, which is not conducive to fatigue performance in steel.
  • the content of P, S, and O in steel needs to be as low as possible, but considering the economics of steel smelting costs, the mass percentage of the above inevitable impurity elements is controlled in a certain appropriate range, when the inevitable When the impurity elements are controlled within the appropriate range, the harmful effects of the inevitable impurity elements can be minimized, so that there is no obvious adverse effect on the performance of the steel.
  • the other inevitable impurities satisfy at least one of the following items: P ⁇ 0.015%, S ⁇ 0.0020%, O ⁇ 0.003% .
  • each related element also satisfies: Ti / N ⁇ 5, which is because by controlling Ti / N to ⁇ 5, it is beneficial to retain sufficient Ti element Reacts with C to form TiC precipitation strengthening.
  • the microstructure is ferrite + lower bainite, and the average width of the lower bainite lath is ⁇ 500 nm. 5-10%.
  • the average width of the lower bainite lath is ⁇ 400 nm.
  • nano-scale TiC interphase precipitates are formed in the ferrite, and the particle diameter of 70% or more of the TiC interphase precipitates in the ferrite is Below 30nm. This is for the present case, the finer the precipitates, the better the precipitation strengthening effect.
  • the grades of all types of non-metallic inclusions are below grade 1.0, and the total rating of all non-metallic inclusions is controlled below 3.0, and it does not have long strips Shaped inclusions.
  • the rating of non-metallic inclusions can be controlled to no greater than 1.0, and the overall rating of all non-metallic inclusions should be controlled below 3.0, while suppressing long strips
  • another object of the present invention is to provide a method for manufacturing the above 800MPa hot stamping axle housing steel.
  • the 800MPa hot stamping axle housing steel obtained by the manufacturing method can reach the strength level of 800MPa, while the 800MPa
  • the plasticity and fatigue performance of the grade hot stamping axle housing steel is also good, which is very suitable for manufacturing axle housing.
  • the present invention proposes a method for manufacturing the above 800MPa hot stamping axle housing steel, which includes the steps of:
  • Cooling cooling in two stages after rolling, first cooling the steel plate at a rate of 80-200 ° C / s to 680-730 ° C, natural air cooling for 5-7s; then cooling the steel plate at a rate of 30-70 ° C / s To 360 ⁇ 450 °C, then coiled or natural air cooled to room temperature.
  • the manufacturing method can obtain a small amount of ferrite + lower bainite structure through sectional cooling control combined with coiling in the middle temperature range of 360 to 450 ° C,
  • the average width of the lower bainite lath is ⁇ 500 nm, so that the steel plate finally obtained has a yield strength ⁇ 800 MPa, a tensile strength ⁇ 900 MPa, an elongation A 50 ⁇ 22%, and an impact energy of -20 ° C above 60J.
  • step (3) the reduction rate of the last pass rolling is controlled to be> 15%; the final rolling temperature is 820 to 900 ° C, in order to enable the austenite to accumulate enough before the bainite transformation Deformation promotes the formation of finer bainite structures. If the rolling temperature is too low, the high-temperature phase transformation of ferrite is likely to occur, which reduces the strength of the steel; if the rolling temperature is too high, the deformation accumulated in the austenite will be restored, which is not conducive to refining the structure after phase transformation.
  • step (4) the steel plate is first cooled at a cooling rate of 80 to 200 ° C / s to 680 to 730 ° C, in order to quickly cool to the temperature range of ferrite formation, natural air cooling is 5-7s, and the formation ratio is 5-10% ferrite.
  • a cooling rate of 80 to 200 ° C / s to 680 to 730 ° C in order to quickly cool to the temperature range of ferrite formation, natural air cooling is 5-7s, and the formation ratio is 5-10% ferrite.
  • the steel plate is rapidly cooled to a lower temperature of 360-450 ° C at a cooling rate of 30-70 ° C / s, and then coiled or naturally air-cooled to room temperature.
  • rapid cooling is to suppress the ferrite from continuing to change phase It is cooled to 360 ⁇ 450 °C coiling or natural air cooling is to form a fine lower bainite structure, while keeping the temperature at a lower temperature is conducive to inhibit the continued growth of precipitate particles.
  • step (2) the slab is heated in a furnace with a target temperature range of 1180 to 1270 ° C. After warming to the target temperature, heat preservation starts, and the heat preservation time is> 1.5h.
  • the slab is heated in a furnace with a target temperature ranging from 1180 to 1270 ° C.
  • the heat preservation is started.
  • the heat preservation time is> 1.5h to ensure that the alloy elements are fully solidified. Dissolve.
  • the slab heating furnace has three stages of preheating, heating and soaking. Whether the current core of the slab is heated to the target temperature is calculated according to the calculation model provided by the heating furnace equipment supplier, which can be calculated The time required for the billet in the three stages of preheating, heating, and soaking to reach the target temperature, and the calculation model was checked with a thermocouple in the early stage of development.
  • the heating temperature exceeds 1270 ° C
  • the austenite grains grow excessively, causing the intergranular bonding strength to be weakened, and cracks are easy to occur during rolling; in addition, the heating temperature exceeding 1270 ° C may easily cause decarburization of the slab surface and the mechanics of the final steel Performance is adversely affected.
  • the 800MPa hot stamping axle housing steel described in the present invention is excellent in strength, impact, fatigue, welding, etc. It controls the inclusions that affect fatigue through rationally optimized alloy design, for example, in controlling rolling, Sectional cooling, and coiling in the middle temperature range or natural air cooling to room temperature, to control the formation of the required ferrite + lower bainite, and is conducive to the formation of nano-scale precipitate particles in ferrite, which ultimately makes the present invention
  • the 800 MPa grade hot stamping axle housing steel described above achieves a high plasticity at the 800 MPa yield strength level, that is, an elongation A 50 ⁇ 22%, and its low temperature impact performance and welding performance are good.
  • the 800 MPa hot stamping axle housing steel described in the present invention specifically controls the components and inclusions that may affect the fatigue performance of the steel plate, it is conducive to improving the fatigue performance of the steel plate and is beneficial to suppressing the Austrian The growth of the austenite and refine the final structure, so as to achieve a higher strength after hot stamping of the steel plate.
  • the technical solution described in the present invention obtains a small amount of ferrite + lower bainite structure by controlling the content of Ti, B and other elements, combined with controlled rolling and controlled cooling, which improves the plasticity and fatigue performance of the steel plate and is very suitable for use in Manufacturing axle housings.
  • FIG. 1 is a metallurgical structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 1 is a metallurgical structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 2 is a metallurgical structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • FIG. 1 is a metallurgical structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • FIG. 3 is a scanning metallographic structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 3 is a scanning metallographic structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 4 is a scanning metallographic structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • FIG. 4 is a scanning metallographic structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • FIG. 5 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 1.
  • FIG. 5 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 1.
  • FIG. 6 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 6 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 7 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 3.
  • FIG. 7 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 3.
  • FIG. 8 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 4.
  • FIG. 8 schematically shows the size, morphology and distribution of typical inclusions of the 800 MPa hot stamping axle housing steel of Example 4.
  • Table 1 lists the mass percentage of each chemical element in the 800 MPa class hot stamping axle housing steel of Examples 1-6.
  • the 800 MPa grade hot stamping axle housing steels of Examples 1-6 were prepared using the following steps (see Table 2 for specific process parameters):
  • the slab is heated in a furnace with a target temperature ranging from 1180 to 1270 ° C. After the core of the slab is heated to the target temperature, heat preservation is started, and the heat preservation time is> 1.5h. ;
  • Cooling cooling in two stages after rolling, first cooling the steel plate at a rate of 80-200 ° C / s to 680-730 ° C, natural air cooling for 5-7s; then cooling the steel plate at a rate of 30-70 ° C / s To 360 ⁇ 450 °C, then coiled or natural air cooled to room temperature.
  • Table 2 lists the specific process parameters of the method for manufacturing the 800 MPa hot stamping axle housing steel of Examples 1-6.
  • the performance test was performed on the 800 MPa grade hot stamping axle housing steel of Examples 1-6, and the test results are listed in Table 3. Among them, the tensile test (yield strength, tensile strength, elongation) is tested using GB / T 228.1-2010 "Metal material room temperature tensile test method” standard, and the impact test GB / T 229-2007 “Metal material Charpy pendulum Impact test method "standard test.
  • the 800 MPa grade hot stamped axle housing steel of Examples 1-6 has a yield strength ⁇ 800 MPa, a tensile strength ⁇ 900 MPa, an elongation A 50 ⁇ 22%, and an impact energy -20 ° C ⁇ 60J. It can be seen from this that the 800 MPa class hot stamping axle housing steel of each embodiment of the present case can reach 800 MPa strength level. At the same time, the 800 MPa class hot stamping axle housing steel has good plasticity and fatigue performance, and is very suitable for manufacturing axle housings.
  • Figure 1 and Figure 2 were taken with LEICACTR6500 optical metallurgical microscope, which was taken at 200 times magnification.
  • Figures 3 and 4 use the JCM7000 scanning electron microscope, which was taken at 20,000 and 10,000 times magnification, respectively.
  • Figures 5 to 8 were taken with the LEICACTR6500 optical metallurgical microscope, magnified 50 times and referenced to the national standard "Microscopic Inspection Method for the Standard Rating Chart for the Determination of the Content of Nonmetallic Inclusions in Steel" (GB / T 10561-2005).
  • FIG. 1 is a metallurgical structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 2 is a metallurgical structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • FIG. 1 is a metallurgical structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 2 is a metallurgical structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • microstructure of the 800 MPa grade hot stamping axle housing steel of Examples 2 and 3 of the present case is ferrite + lower bainite, and the comparative example of ferrite is 5-10%.
  • FIG. 3 is a scanning metallographic structure diagram of 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 4 is a scanning metallographic structure diagram of 800 MPa grade hot stamping axle housing steel of Example 3.
  • the average width of the lower bainite lath in FIG. 4 is ⁇ 400 nm.
  • Nanoscale TiC interphase precipitates are formed in the ferrite of 800 MPa grade hot stamped axle housing steel in Example 2 of this case, where the particle diameter of more than 70% of the TiC interphase precipitates in ferrite is below 30 nm.
  • the 800 MPa hot stamping axle housing steel of Example 3 lower bainite is precipitated, and the width of the lower bainite lath is 300 nm or less.
  • FIGS. 5 to 8 and Table 4 schematically shows the size, morphology and distribution of inclusions of the 800 MPa hot stamping axle housing steel of Example 1.
  • 6 schematically shows the size, morphology and distribution of inclusions in the 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 7 schematically shows the size, morphology and distribution of inclusions in the 800 MPa hot stamping axle housing steel of Example 3.
  • FIG. 8 schematically shows the size, morphology and distribution of inclusions in the 800 MPa hot stamping axle housing steel of Example 4.
  • FIG. 5 schematically shows the size, morphology and distribution of inclusions of the 800 MPa hot stamping axle housing steel of Example 1.
  • 6 schematically shows the size, morphology and distribution of inclusions in the 800 MPa hot stamping axle housing steel of Example 2.
  • FIG. 7 schematically shows the size, morphology and distribution of inclusions in the 800 MPa hot stamping axle housing steel of Example 3.
  • FIG. 8 schematically shows the size, morphology and distribution of inclusions in the 800 MP
  • Table 4 lists the non-metallic inclusion rating results for the 800 MPa hot stamping axle housing steel of Examples 1-6.
  • the 800 MPa hot stamping axle housing steel described in this case is designed by reasonably optimizing alloy elements, and controlling the ratio of each alloy element and the level of inclusions, and at the same time, the process is controlled by rolling and cooling to obtain The required microstructure is obtained, a small amount of ferrite + lower bainite is obtained, and a large number of nano-scale TiC particles are formed in the ferrite, so that the final 800MPa hot stamping bridge shell steel can reach more than 800MPa
  • the yield strength, combined with good plasticity, low temperature toughness and fatigue performance, is very suitable for high-strength weight reduction of hot stamping axle housings.

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Abstract

一种800MPa级热冲压桥壳钢，其化学元素质量百分比为：C：0.15-0.21%，Si：0.30-0.80%，Mn：1.75-2.10%，Nb：0.015-0.040%，Ti：0.020-0.060%，B：0.0015-0.0030%，Al：0.005-0.015%，Ca：0.0004-0.001%，N：0.001-0.001%，余量为Fe及其他不可避免的杂质。此外还提供了制备上述800MPa级热冲压桥壳钢的制造方法，包括以下步骤：（1）冶炼和铸造；（2）加热；（3）轧制：控制最后一道次轧制的压下率＞15%；终轧温度为820-900℃；（4）冷却：轧后分两段冷却。

Description

一种800MPa级热冲压桥壳钢及其制造方法 技术领域
本发明涉及一种钢种及其制造方法，尤其涉及一种800MPa级热冲压桥壳钢及其制造方法。
背景技术
桥壳作为汽车关键承载部件，对安全性有较高要求，需要满足严格的构件疲劳性能，也就是说对于桥壳用钢而言，要求钢板性能稳定、耐低温冲击、焊接性良好，同时对偏析和夹杂物进行严格控制。
目前，我国热冲压桥壳钢主要为16Mn、Q345C、Q420C、Q460C等普通C-Mn结构钢，此类C-Mn钢热冲压后强度进一步下降，例如Q460C热冲压后屈服强度下降到400MPa左右。
公开号为CN104213019A，公开日为2014年12月17日，名称为“一种600MPa级汽车桥壳钢及其生产方法”的中国专利文献公开了一种600MPa级汽车桥壳钢及其生产方法。在该专利文献所公开的技术方案中，其通过控制V和N元素的含量的同时，控轧控冷，最终获得600MPa级汽车桥壳用热轧带钢。
公开号为CN103422020A，公开日为2013年12月4日，名称为“一种冲焊桥壳用钢板及其制造方法”的中国专利文献公开了一种冲焊桥壳用钢板及其制造方法。在该专利文献所公开的技术方案中，其综合通过添加Nb、V等元素提高热冲压后的钢板强度，并通过对Ti、Al等元素的综合控制，提高钢板热冲压后的低温韧性和疲劳性能，最终得到600MPa级强度钢板。
基于此，期望获得一种强度较高的桥壳钢，其可以达到800MPa强度级别，同时钢板的塑性和疲劳性能也表现良好，从而更适合用于制造车桥壳。
发明内容
本发明的目的之一在于提供一种800MPa级热冲压桥壳钢，其可以达到800MPa强度级别，同时该800MPa级热冲压桥壳钢的塑性和疲劳性能也表现 良好，非常适合用于制造车桥壳。
为了实现上述目的，本发明提出了一种800MPa级热冲压桥壳钢，其化学元素质量百分比为：
C：0.15～0.21％，Si：0.30～0.80％，Mn：1.75～2.10％，Nb：0.015～0.040％，Ti：0.020～0.060％，B：0.0015～0.0030％，Al：0.005～0.015％，Ca：0.0004～0.001％，N：0.001～0.004％，余量为Fe及其他不可避免的杂质。
本发明所述的800MPa级热冲压桥壳钢中的各化学元素的设计原理如下所述：
C：在本发明所述的800MPa级热冲压桥壳钢中，碳起到固溶强化作用，添加C可以提高下贝氏体的强度，此外，C还可以在桥壳热冲压空冷过程中与Fe发生反应，形成弥散分布的Fe 3C，从而提高热冲压后钢板的强度。但是C的质量百分比过高，则不利于钢板的焊接性。因此，在本发明所述的800MPa级热冲压桥壳钢中，控制C的质量百分比在0.15～0.21％。
Si：在本发明所述的800MPa级热冲压桥壳钢中，Si可以抑制渗碳体在高温下的析出，从而有利于下贝氏体的形成，同时Si还可以使得钢板中形成细小的渗碳体颗粒，以提高下贝氏体的强度。然而，过高含量的Si则不利于钢板的焊接性能。因此，在本发明所述的技术方案中控制Si的质量百分比在0.30～0.80％。
Mn：对于本发明所述的技术方案中，由于添加一定含量Mn元素有利于促进形成下贝氏体，并且Mn可以对下贝氏体组织起到一定的固溶强化作用。此外，添加Mn元素还有利于在钢板热冲压时形成较细的铁素体或贝氏体，从而有利于提高热冲压后钢板的强度。然而，过高含量的Mn不利于钢板的焊接性能。因此，在本发明所述的800MPa级热冲压桥壳钢中Mn的质量百分比为1.75～2.10％。
Nb：对于本发明所述的800MPa级热冲压桥壳钢而言，其在控轧控冷阶段阶段，钢板中的微量的Nb与C反应形成细小的NbC颗粒，有利于细化钢板组织，同时还可以提高钢板的强度、塑性和韧性。此外，NbC在下贝氏体组织中析出，可以产生较强的析出强化作用。另外，在热冲压加热阶段NbC还可以细化奥氏体晶粒，提高桥壳钢板热冲压后的强度。然而，Nb元素含量太高时，NbC颗粒较大，抑制奥氏体晶粒长大的效果反而弱化。因此，在本发明所 述的800MPa级热冲压桥壳钢中Nb的质量百分比控制在0.015～0.040％。
Ti：在本发明所述的技术方案中，Ti元素在钢板轧制后奥氏体向铁素体相变过程中会与C反应形成直径几纳米到几十纳米的TiC颗粒，从而产生析出强化作用，尤其是在680～730℃范围内可以产生更细的相间析出。此外，在热冲压加热阶段TiC细小颗粒可以抑制奥氏体晶粒的长大，进而细化热冲压后的组织，提高桥壳钢板热冲压后的强度。然而Ti元素含量太高时，容易与N反应形成微米级立方体型TiN大颗粒，恶化钢板的韧性和疲劳性能。因此，在本发明所述的800MPa级热冲压桥壳钢中Ti的质量百分比控制在0.020～0.060％。
B：在本发明所述的800MPa级热冲压桥壳钢中，微量的B可以促进贝氏体的形成，然而，B元素含量太高时容易产生B脆问题，恶化钢板的冲击韧性。此外，在桥壳热冲压阶段，微量的B元素有利于促进形成较细的贝氏体，提高钢板的强度。因此，在本发明所述的800MPa级热冲压桥壳钢中B的质量百分比控制在0.0015～0.0030％。
Al：对于本发明所述的800MPa级热冲压桥壳钢而言，Al作为重要的脱氧剂，通常加入0.02％以上的Al。然而，考虑到桥壳钢对疲劳性能的严格要求，需要对Al的氧化物链状夹杂进行控制，因此，将Al含量控制在较低的范围。因此，在本发明所述的800MPa级热冲压桥壳钢中Al的质量百分比控制在0.005～0.015％。
Ca：对于本发明所述的800MPa级热冲压桥壳钢而言，微量的Ca元素可以在钢冶炼过程中的起到净化剂作用，改善钢的韧性和疲劳性能；同时，Ca处理可以改善MnS夹杂的形状，防止形成长条形的MnS夹杂物。然而，Ca含量超过0.001％容易形成尺寸较大的Ca的化合物，反而会恶化韧性和疲劳性能。因此，在本发明所述的800MPa级热冲压桥壳钢中Ca的质量百分比控制在0.0004～0.001％。
N：对于本发明所述的800MPa级热冲压桥壳钢而言，本案需要将N元素控制在较窄的范围。微量的N元素可以与Ti反应形成TiN颗粒，在焊接和热冲压时，可以有效抑制奥氏体晶粒的长大，细化焊接热影响区和热冲压后的组织，提高热影响区和热冲压钢板的强度、低温韧性和疲劳性能。然而，N含量太高时，形成的TiN颗粒太大，反而会恶化钢板低温韧性和疲劳性能。因此，在本发明所述的800MPa级热冲压桥壳钢中N的质量百分比控制在 0.001～0.004％。
需要说明的是，在本发明所述的技术方案中，其他不可避免的杂质元素包括P、S和O。P、S和O都属于钢材料中不可避免的有害杂质元素，不利于钢性能，例如P作为杂质元素容易产生冷脆问题；S容易与Mn反应产生MnS夹杂，不利于钢中疲劳性能；O容易与Al反应产生Al 3O链状夹杂，不利于钢中疲劳性能。因此，在钢中P、S以及O的含量需要越低越好，但是考虑到钢铁冶炼成本的经济性，因而，对于上述不可避免的杂质元素的质量百分比控制在一定的适当范围，当不可避免的杂质元素控制在该适当范围内时，可以将不可避免的杂质元素的有害作用降至最低，从而不会对钢性能产生明显不利影响。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，所述其他不可避免的杂质满足下述各项的至少其中之一：P≤0.015％，S≤0.0020％，O≤0.003％。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，各相关元素还满足：Ti/N≥5，这是由于通过将Ti/N控制在≥5时，有利于保留足够Ti元素与C发生反应形成TiC析出强化。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，其微观组织为铁素体+下贝氏体，下贝氏体板条的平均宽度≤500nm，其中铁素体的相比例为5-10％。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，下贝氏体板条的平均宽度≤400nm。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，所述铁素体中形成有纳米级的TiC相间析出物，铁素体中70％以上的TiC相间析出物的颗粒直径在30nm以下。这是对于本案而言，析出物的颗粒越细，则析出强化效果越好。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，其屈服强度≥800MPa，抗拉强度≥900MPa，延伸率A 50≥22％，-20℃冲击功≥60J。
进一步地，在本发明所述的800MPa级热冲压桥壳钢中，其各类非金属夹杂物的等级在1.0级以下，所有非金属夹杂物评级总和控制在3.0以下，并且其不具有长条形夹杂物。
在上述方案中，考虑到桥壳钢对疲劳性能的要求，因而，可以将非金属夹 杂物评级控制在不大于1.0级，所有非金属夹杂物评级综合控制在3.0级以下，同时抑制长条形夹杂物的形成，以提高本案所述的800MPa级热冲压桥壳钢的性能。
相应地，本发明的另一目的在于提供一种上述的800MPa级热冲压桥壳钢的制造方法，通过该制造方法获得的800MPa级热冲压桥壳钢，其可以达到800MPa强度级别，同时该800MPa级热冲压桥壳钢的塑性和疲劳性能也表现良好，非常适合用于制造车桥壳。
为了实现上述目的，本发明提出了一种上述的800MPa级热冲压桥壳钢的制造方法，其包括步骤：
(1)冶炼和铸造；
(2)加热；
(3)轧制：控制最后一道次轧制的压下率>15％；终轧温度为820～900℃；
(4)冷却：轧后分两段冷却，首先将钢板以80～200℃/s的速度冷却至680～730℃，自然空冷5-7s；然后将钢板以30～70℃/s的速度冷却至360～450℃，然后卷取或自然空冷至室温。
在本发明所述的800MPa级热冲压桥壳钢的制造方法中，所述制造方法通过分段冷却控制配合360～450℃中温范围卷取可以得到少量的铁素体+下贝氏体组织，下贝氏体板条的平均宽度≤500nm，以使得最终获得的钢板屈服强度≥800MPa，抗拉强度≥900MPa，延伸率A 50≥22％，-20℃冲击功在60J以上。
其中，在步骤(3)中将控制最后一道次轧制的压下率>15％；终轧温度为820～900℃，是为了可以使得奥氏体往贝氏体相变前积累足够多的变形，促进形成较细的贝氏体组织。轧制温度太低，容易发生铁素体的高温相变，降低钢的强度；轧制温度太高，奥氏体中积累的变形发生回复，不利于细化相变后的组织。
而在步骤(4)中，首先将钢板以80～200℃/s的冷速冷至680～730℃，是为了快速冷却到铁素体形成温度区间，自然空冷5-7s，形成相比例为5-10％的铁素体。在680-730℃的高温下，使得奥氏体向铁素体相变时会产生细密的TiC颗粒相间析出。然后将钢板以30～70℃/s的冷速快速冷至360～450℃的较低温度，然后卷取或自然空冷至室温，在上述步骤中，快速冷却是为了抑制铁素体继续相变，冷却至360～450℃卷取或自然空冷是为了形成细密的下贝氏体组织， 同时在较低的温度保温有利于抑制析出物颗粒继续长大。
进一步地，在本发明所述的800MPa级热冲压桥壳钢的制造方法中，在步骤(2)中：将铸坯在目标温度范围为1180～1270℃的炉中加热，待铸坯心部升温至目标温度后开始保温，保温时间>1.5h。
上述方案中，将将铸坯在目标温度范围为1180～1270℃的炉中加热，待铸坯心部升温至某一目标温度后开始保温，保温时间>1.5h，是为了保证合金元素充分固溶。其中，板坯加热炉有预热、加热、均热三段，当前的铸坯心部是否升温至目标温度，是根据加热炉设备供应商提供的计算模型进行计算的，该计算模型可计算出处于预热、加热、均热三段中的铸坯达到目标温度所需的时间，而且该计算模型在开发初期采用热电偶进行校核。而加热温度超过1270℃时奥氏体晶粒过度长大，引起晶间结合力减弱，在轧制时容易产生裂纹；此外，加热温度超过1270℃容易引起钢坯表面脱碳，对最终钢的力学性能造成不良影响。
本发明所述的800MPa级热冲压桥壳钢及其制造方法的优点及有益效果如下所述：
本发明所述的800MPa级热冲压桥壳钢在强度、冲击、疲劳、焊接等方面表现均较为优异，其通过合理优化的合金设计，对影响疲劳的夹杂物进行控制，例如在控制轧制，分段冷却，以及在中温范围卷取或自然空冷至室温，以控制形成需要的铁素体+下贝氏体，并且有利于铁素体中形成纳米级的析出物颗粒，最终使得本发明所述的800MPa级热冲压桥壳钢在实现800MPa屈服强度级别的同时，实现了较高的塑性，即延伸率A 50≥22％，并且其低温冲击性能和焊接性能良好。
另外，由于本发明所述的800MPa级热冲压桥壳钢对可能影响钢板疲劳性能的成分和夹杂物进行专门控制，有利于提高钢板的疲劳性能，在进行桥壳热冲压过程中有利于抑制奥氏体的长大并细化最终组织，从而实现钢板热冲压后具有较高的强度。
此外，本发明所述的技术方案通过控制Ti、B等元素的含量，配合控轧控冷，得到少量的铁素体+下贝氏体组织，提高钢板的塑性和疲劳性能，十分适合用于制造车桥壳。
附图说明
图1为实施例2的800MPa级热冲压桥壳钢的金相组织图。
图2为实施例3的800MPa级热冲压桥壳钢的金相组织图。
图3为实施例2的800MPa级热冲压桥壳钢的扫描金相组织图。
图4为实施例3的800MPa级热冲压桥壳钢的扫描金相组织图。
图5示意性显示实施例1的800MPa级热冲压桥壳钢的典型夹杂物尺寸、形貌和分布。
图6示意性显示实施例2的800MPa级热冲压桥壳钢的典型夹杂物尺寸、形貌和分布。
图7示意性显示实施例3的800MPa级热冲压桥壳钢的典型夹杂物尺寸、形貌和分布。
图8示意性显示实施例4的800MPa级热冲压桥壳钢的典型夹杂物尺寸、形貌和分布。
具体实施方式
下面将结合说明书附图和具体的实施例对本发明所述的800MPa级热冲压桥壳钢及其制造方法做进一步的解释和说明，然而该解释和说明并不对本发明的技术方案构成不当限定。
实施例1-6
表1列出了实施例1-6的800MPa级热冲压桥壳钢中各化学元素的质量百分比。
表1.(wt％，余量为Fe和除P、S以及O以外的不可避免的杂质)
Figure PCTCN2019111982-appb-000001
Figure PCTCN2019111982-appb-000002
实施例1-6的800MPa级热冲压桥壳钢采用下述步骤制得(具体工艺参数参见表2)：
(1)冶炼和铸造：根据表1所列出的化学元素的质量百分比进行配比冶炼，采用真空电路进行冶炼，随后将冶炼的钢水浇注成铸坯；
(2)加热：将铸坯在目标温度范围为1180～1270℃的炉中加热，待铸坯心部升温至目标温度后开始保温，保温时间>1.5h。；
(3)轧制：控制最后一道次轧制的压下率>15％；终轧温度为820～900℃
(4)冷却：轧后分两段冷却，首先将钢板以80～200℃/s的速度冷却至680～730℃，自然空冷5-7s；然后将钢板以30～70℃/s的速度冷却至360～450℃，然后卷取或自然空冷至室温。
表2列出了实施例1-6的800MPa级热冲压桥壳钢的制造方法的具体工艺参数。
表2.
Figure PCTCN2019111982-appb-000003
对实施例1-6的800MPa级热冲压桥壳钢进行了性能测试，测试结果列于表3中。其中拉伸试验(屈服强度、抗拉强度、延伸率)采用GB/T 228.1-2010《金属材料室温拉伸试验方法》标准进行测试，冲击试验GB/T 229-2007《金属材料夏比摆锤冲击试验方法》标准测试。
表3.
Figure PCTCN2019111982-appb-000004
注：表3中的-20℃冲击功试验结果中三列分别代表三个平行试样的测试结果
由表3可以看出，实施例1-6的800MPa级热冲压桥壳钢，其屈服强度≥800MPa，抗拉强度≥900MPa，延伸率A 50≥22％，-20℃冲击功≥60J。由此说明，本案各实施例的800MPa级热冲压桥壳钢可以达到800MPa强度级别，同时该800MPa级热冲压桥壳钢的塑性和疲劳性能也表现良好，非常适合用于制造车桥壳。
其中图1和图2采用LEICACTR6500光学金相显微镜，放大200倍拍摄。图3和图4采用JCM7000扫描电镜，分别放大20000和10000倍拍摄。图5-图8采用LEICACTR6500光学金相显微镜，放大50倍并参照国家标准《钢中非金属夹杂物含量的测定标准评级图显微检验法》(GB/T 10561-2005)拍摄。
图1为实施例2的800MPa级热冲压桥壳钢的金相组织图。图2为实施例3的800MPa级热冲压桥壳钢的金相组织图。
结合图1和图2可以看出，本案实施例2和3的800MPa级热冲压桥壳钢的微观组织为铁素体+下贝氏体，其中铁素体的相比例为5-10％。
图3为实施例2的800MPa级热冲压桥壳钢的扫描金相组织图。图4为实施例3的800MPa级热冲压桥壳钢的扫描金相组织图。
结合图3和图4可以看出，图4中下贝氏体板条的平均宽度≤400nm。本案实施例2的800MPa级热冲压桥壳钢的铁素体中形成有纳米级的TiC相间析出物，其中铁素体中70％以上的TiC相间析出物的颗粒直径在30nm以下，而在实施例3的800MPa级热冲压桥壳钢中析出有下贝氏体，下贝氏体板条宽度在300nm以下。
此外，由于夹杂物对于钢板性能具有一定影响，因而，对本案实施例1-4的夹杂物进行了测试，所得到的结果显示于图5至图8以及表4中。其中，图， 5示意性显示实施例1的800MPa级热冲压桥壳钢的夹杂物尺寸、形貌和分布。图6示意性显示实施例2的800MPa级热冲压桥壳钢的夹杂物尺寸、形貌和分布。图7示意性显示实施例3的800MPa级热冲压桥壳钢的夹杂物尺寸、形貌和分布。图8示意性显示实施例4的800MPa级热冲压桥壳钢的夹杂物尺寸、形貌和分布。
表4列出了实施例1-6的800MPa级热冲压桥壳钢的非金属夹杂物评级结果。各类夹杂物的定义和评级方法参照国家标准《钢中非金属夹杂物含量的测定标准评级图显微检验法》(GB/T 10561-2005)。
表4.
Figure PCTCN2019111982-appb-000005
结合图5至图8以及表4可以看出，在本案实施例1-6的800MPa级热冲压桥壳钢中，其各类非金属夹杂物的等级在1.0级以下，所有非金属夹杂物评级总和控制在3.0以下，并且其不具有长条形夹杂物。
综上所述可以看出，本案所述的800MPa级热冲压桥壳钢通过合理优化的合金元素设计，并控制各合金元素间的比例以及夹杂物水平，同时配合工艺控轧控冷，从而获得了所需要的微观组织结构，得到了少量的铁素体+下贝氏体，同时在铁素体中形成大量的纳米级TiC颗粒，以使得最终的800MPa级热冲压桥壳钢可以达到800MPa以上的屈服强度，并兼具有良好的塑性、低温韧性和疲劳性能，十分适合热冲压桥壳的高强减重使用。
需要说明的是，本发明的保护范围中现有技术部分并不局限于本申请文件所给出的实施例，所有不与本发明的方案相矛盾的现有技术，包括但不局限于在先专利文献、在先公开出版物，在先公开使用等等，都可纳入本发明的保护 范围。
此外，本案中各技术特征的组合方式并不限本案权利要求中所记载的组合方式或是具体实施例所记载的组合方式，本案记载的所有技术特征可以以任何方式进行自由组合或结合，除非相互之间产生矛盾。
还需要注意的是，以上所列举的实施例仅为本发明的具体实施例。显然本发明不局限于以上实施例，随之做出的类似变化或变形是本领域技术人员能从本发明公开的内容直接得出或者很容易便联想到的，均应属于本发明的保护范围。

Claims (10)

  1. 一种800MPa级热冲压桥壳钢，其特征在于，其化学元素质量百分比为：
    C：0.15～0.21％，Si：0.30～0.80％，Mn：1.75～2.10％，Nb：0.015～0.040％，Ti：0.020～0.060％，B：0.0015～0.0030％，Al：0.005～0.015％，Ca：0.0004～0.001％，N：0.001～0.004％，余量为Fe及其他不可避免的杂质。
  2. 如权利要求1所述的800MPa级热冲压桥壳钢，其特征在于，所述其他不可避免的杂质满足下述各项的至少其中之一：P≤0.015％，S≤0.0020％，O≤0.003％。
  3. 如权利要求1所述的800MPa级热冲压桥壳钢，其特征在于，各相关元素还满足：Ti/N≥5。
  4. 如权利要求1所述的800MPa级热冲压桥壳钢，其特征在于，其微观组织为铁素体+下贝氏体，所述下贝氏体板条的平均宽度≤500nm，其中铁素体的相比例为5-10％。
  5. 如权利要求4所述的800MPa级热冲压桥壳钢，其特征在于，所述下贝氏体板条的平均宽度≤400nm。
  6. 如权利要求4所述的800MPa级热冲压桥壳钢，其特征在于，所述铁素体中形成有纳米级的TiC相间析出物，铁素体中70％以上的TiC相间析出物的颗粒直径在30nm以下。
  7. 如权利要求1所述的800MPa级热冲压桥壳钢，其特征在于，其屈服强度≥800MPa，抗拉强度≥900MPa，延伸率A 50≥22％，-20℃冲击功≥60J。
  8. 如权利要求1所述的800MPa级热冲压桥壳钢，其特征在于，其各类非金属夹杂物的等级在1.0级以下，所有非金属夹杂物评级总和控制在3.0以下，并且其不具有长条形夹杂物。
  9. 如权利要求1-8中任意一项所述的800MPa级热冲压桥壳钢的制造方法，其包括步骤：
    (5)冶炼和铸造；
    (6)加热；
    (7)轧制：控制最后一道次轧制的压下率>15％；终轧温度为820～900℃；
    (8)冷却：轧后分两段冷却，首先将钢板以80～200℃/s的速度冷却至680～730℃，自然空冷5-7s；然后将钢板以30～70℃/s的速度冷却至360～450℃，然后卷取或自然空冷至室温。
  10. 如权利要求9所述的800MPa级热冲压桥壳钢的制造方法，其特征在于，在步骤(2)中：将铸坯在目标温度范围为1180～1270℃的炉中加热，待铸坯心部升温至目标温度后开始保温，保温时间>1.5h。
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CN113528779A (zh) * 2021-06-22 2021-10-22 首钢集团有限公司 一种控制桥壳钢热成形后强度下降的方法及制得的桥壳钢
CN114032455A (zh) * 2021-10-18 2022-02-11 武汉钢铁有限公司 一种易焊接800MPa级热连轧桥壳钢及生产方法
CN114032455B (zh) * 2021-10-18 2022-08-26 武汉钢铁有限公司 一种易焊接800MPa级热连轧桥壳钢及生产方法
CN114645122A (zh) * 2022-02-16 2022-06-21 包头钢铁（集团）有限责任公司 一种稀土处理的高强高韧桥壳钢板的制备方法
CN114645122B (zh) * 2022-02-16 2024-06-11 包头钢铁（集团）有限责任公司 一种稀土处理的高强高韧桥壳钢板的制备方法
CN114561505A (zh) * 2022-03-22 2022-05-31 承德建龙特殊钢有限公司 一种低温环境车桥管用连铸圆坯及其制备方法
CN118422058A (zh) * 2024-05-02 2024-08-02 新疆八一钢铁股份有限公司 一种高韧性Q420MPa风电用钢板的冶炼方法

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