WO2020143367A1 - 一种690MPa级特厚钢板及其制造方法 - Google Patents

一种690MPa级特厚钢板及其制造方法 Download PDF

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WO2020143367A1
WO2020143367A1 PCT/CN2019/122904 CN2019122904W WO2020143367A1 WO 2020143367 A1 WO2020143367 A1 WO 2020143367A1 CN 2019122904 W CN2019122904 W CN 2019122904W WO 2020143367 A1 WO2020143367 A1 WO 2020143367A1
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steel plate
slab
temperature
thick steel
extra
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PCT/CN2019/122904
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French (fr)
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孙超
李东晖
尹雨群
赵柏杰
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南京钢铁股份有限公司
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Priority to JP2021538959A priority Critical patent/JP7267430B2/ja
Publication of WO2020143367A1 publication Critical patent/WO2020143367A1/zh

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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • 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
    • C21D9/0081—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for slabs; for billets
    • 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
    • C22C38/18—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
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese

Definitions

  • the invention relates to a special thick steel plate and a manufacturing method thereof, in particular to a 690MPa class special thick steel plate and a manufacturing method thereof.
  • the Chinese invention patent with the patent number 201510125485.1 discloses a low-yield-ratio high-strength thick steel plate with excellent low-temperature impact toughness and a manufacturing method thereof.
  • the chemical composition of the low-yield-ratio high-strength thick steel plate contains 3.6-5.5% Ni is expensive.
  • the Chinese patent application with the application number 201610026446.0 discloses a high-strength steel plate for marine engineering and its production method, which adopts the Nb and V microalloying method, but the content of the element is not high due to the hardenability improvement of Mn and the maximum thickness of the steel plate Failed to exceed 100mm.
  • the present invention provides a 690 MPa class extra-thick steel plate, which has excellent mechanical properties of the heart, and can meet the demand for high-performance extra-thick steel plates in harsh service environments such as marine engineering.
  • Another object of the present invention is to provide a method for manufacturing the above-mentioned 690 MPa class extra-thick steel plate.
  • a 690MPa grade extra-thick steel plate according to the present invention the mass percentage content of chemical components are: C: 0.04-0.08%, Mn: 5.2-6.0%, Si: 0.1-0.4%, Mo: 0.1-0.5 %, Ni: 0.2-0.6%, Cr: 0.2-0.6%, Ti: 0.01-0.05%, S: ⁇ 0.005%, P: ⁇ 0.010%, and the balance of Fe and impurities.
  • the thickness of the steel plate is 80-150 mm.
  • microstructure of the steel plate has martensite and austenite; wherein, the volume fraction of austenite is 4-10%.
  • the mechanical properties of the core of the steel plate are: the yield strength is not less than 690MPa, the tensile strength is not less than 770MPa, the elongation after break is not less than 14%, and the energy absorption of the V-shaped sample at -60°C Charpy pendulum impact test Not less than 80J.
  • the steel sheet has a cross-sectional shrinkage rate of not less than 50% based on stretching in the thickness direction.
  • the method for manufacturing a 690 MPa grade extra-thick steel plate according to the present invention adopts the following technical steps:
  • converter smelting is carried out to reduce the S and P contents in molten steel to S ⁇ 0.005%, P ⁇ 0.010%;
  • Tempering heat treatment heating the quenched steel plate to 610-640°C, soaking time 40-70min, after tempering the steel plate is air cooled to room temperature.
  • the steel plate of the present invention uses manganese as the main alloy element, which reduces the amount of nickel and other expensive elements added during the manufacture of extra-thick steel plates, thereby reducing the alloy cost. Adopting a specific manufacturing process, the ultra-thick steel plate with excellent strength, high plasticity and high toughness has excellent mechanical properties and resistance to lamellar tearing, which can meet the demand for high performance in harsh service environments such as marine engineering. The demand for thick steel plates.
  • FIG. 1 is a transmission electron micrograph of the core structure of the steel plate in Example 1.
  • FIG. 1 is a transmission electron micrograph of the core structure of the steel plate in Example 1.
  • C can significantly improve the strength of the structure through interstitial solid solution strengthening. It is an important strengthening element and an important austenite stabilizing element, but in order to ensure low temperature impact toughness And weldability, it is necessary to control the amount of addition at a low level.
  • Mn can improve the strength of the structure through replacement solid solution strengthening, and can also significantly improve the stability of austenite. Appropriate addition of C and Mn can significantly improve the hardenability, reduce the transformation temperature of supercooled austenite, and obtain a high-strength martensite structure.
  • C and Mn reduces the temperature required to form a certain amount of reverse-transformed austenite.
  • C and Mn are enriched in reverse-transformed austenite, so that they can still maintain structural stability at low temperatures, and become an important organization to improve the plasticity and toughness of the present invention.
  • the tempering temperature in the present invention is lower than 600°C, especially when tempering around 550°C, the grain boundaries of elements such as Mn and P are likely to segregate and reduce toughness.
  • the inventor fully considered the mechanism of action of the elements C and Mn in the present invention, and determined the design of the component "manganese in low carbon" with C being 0.04-0.08% and Mn being 5.2-6.0%.
  • Si is a deoxidizing element in the steelmaking process.
  • a proper amount of Si can suppress the segregation of Mn and P and improve toughness.
  • Si can also produce solid solution strengthening, but when the content exceeds 0.3%, the toughness is significantly reduced.
  • the present invention controls Si to 0.1-0.4%.
  • Mo can improve the strength of martensite after tempering, and can also reduce the grain boundary segregation of Mn within a certain content range to improve toughness.
  • the Mo content is controlled at 0.1-0.5%, and the cost of Mo is not significantly increased while playing the role of Mo.
  • Ni can stabilize the austenite phase, improve the hardenability, and reduce the ductile-brittle transition temperature. It is an effective element to improve the low temperature toughness, and it is also conducive to improving the weldability. However, Ni is expensive, and the present invention controls the Ni content to 0.2-0.6%, and fully exerts the beneficial effects of the Ni element without significantly increasing the cost.
  • the Cr can produce a significant solid solution strengthening effect, which is beneficial to increase the strength and improve the corrosion resistance.
  • the Cr content is controlled within a suitable range of 0.2-0.6%.
  • the addition of a small amount of Ti in the present invention can hinder the migration of grain boundaries at a high temperature through a fine and dispersed second phase precipitation form, thereby refining crystal grains and improving mechanical properties.
  • the addition amount is controlled within the range of 0.01-0.05%.
  • the remainder of the present invention is Fe, however, impurities are inevitably introduced from the raw materials or the surrounding environment during the usual manufacturing process. Since these impurities are obvious to those skilled in the art, their names and contents are not specifically described in this specification.
  • the converter smelting is carried out after hot metal desulfurization treatment to reduce the content of S and P to S ⁇ 0.005%, P ⁇ 0.010%, and through a sufficiently high vacuum degree (vacuum degree ⁇ 4mbar) and a sufficiently long vacuum time ( RH treatment (treatment time ⁇ 20min) to reduce the content of gas impurity elements, and the addition of C, Mn, Si, Mo, Ni, Ti and other alloys are completed by LF refining, so it can achieve a high purity smelting effect.
  • a sufficiently high vacuum degree vacuum degree ⁇ 4mbar
  • RH treatment treatment time ⁇ 20min
  • Slab casting can use continuous casting or die casting + forging to obtain slabs of different sizes and specifications.
  • the slab thickness ⁇ 320mm can be obtained by continuous casting, with high production efficiency; the slab with a larger thickness (>320mm) can be die-casted + forged.
  • sufficient rolling total deformation is required as a necessary condition.
  • the ratio of the thickness of the slab to the thickness of the steel plate is required to be ⁇ 4, and the total rolling line rolling variable can be guaranteed to be ⁇ 75%.
  • the thickness of the steel plate is 80mm, the required slab thickness ⁇ 320mm; when the thickness of the steel plate is 150mm, the required slab thickness ⁇ 600mm.
  • the obtained slab obtains the required structure and properties through rolling and heat treatment processes.
  • the Ac3 temperature of the steel is not higher than 770°C.
  • the slab is heated to 1060-1140°C, a high-temperature austenite structure is formed.
  • alloying elements such as C and Mn are homogenized by diffusion.
  • the entire slab achieves homogenization of austenite, and the soaking time of 40-90min can ensure uniform diffusion of elements.
  • the second phase particles of Ti can play a role in inhibiting the growth of crystal grains.
  • the temperature is lower than 1060°C, the element diffusion is too slow, and the austenite homogenization efficiency is too low.
  • the heated slab is recrystallized and rolled at 930°C and above to refine the crystal grains.
  • Open rolling temperature ⁇ 1020°C can avoid excessive growth rate of crystal grains after recrystallization.
  • the pass deformation of ⁇ 10% can make the austenite have enough distortion energy accumulation after deformation to ensure the effect of recrystallization refinement.
  • the critical cooling rate of martensite transformation is lower than 1°C/s, and the martensite structure can be obtained even when the cooling rate is low.
  • the core cooling is usually significantly slower than the surface cooling.
  • the composition design of the present invention can ensure that the martensite transformation also occurs in the core of the steel plate with a thickness of 80-150 mm.
  • the present invention controls the average cooling rate to be not higher than 5°C/s.
  • the redness temperature of the surface of the steel plate cooled after rolling is selected to be 360°C or below, which can avoid obvious element segregation during the cooling process and suppress the precipitation of coarse carbides, and this temperature is also lower than the horse under the composition of the present invention.
  • the cooling process after rolling selected in the present invention can provide a suitable precursor structure for the subsequent heat treatment process.
  • the invention performs quenching + tempering heat treatment on the steel plate.
  • the quenching temperature of 780-830°C is higher than Ac3, soaking to obtain austenite structure. Since the steel plate is water-cooled after rolling, the element segregation and the formation of coarse carbides during the cooling process are avoided, so the element homogenization time in austenite is greatly shortened.
  • the soaking time for quenching and heating is selected to be 5-15 min, which can effectively refine the grain size while ensuring the homogenization of austenite, which is beneficial to the improvement of the mechanical properties of the steel plate.
  • the choice of quenching cooling rate is the same as the reason for selecting the cooling rate after rolling, controlled at 2-8°C/s.
  • the surface redness temperature of the quenched steel sheet is required to be 110° C. or lower, which is lower than the end temperature of martensite transformation under the components of the present invention, and it is possible to ensure that the entire steel sheet obtains a quenched martensite structure with high strength.
  • the tempered heat treatment is performed on the quenched steel plate.
  • reverse transformation austenite with a volume fraction of 4-10% can also be obtained, mainly It is in the form of a thin film, distributed between the martensite laths.
  • austenite stabilizing elements such as C and Mn are enriched in austenite, which improves the stability of austenite.
  • Austenite can still maintain the stability of the crystal structure without martensite transformation.
  • the steel plate is air-cooled, which can also reduce the thermal stress of the extra-thick steel plate and improve the quality of the steel plate.
  • a structure of martensite + austenite is obtained in the entire thickness direction of the steel sheet, particularly in the core of the steel sheet.
  • the martensite matrix provides a yield strength of 690 MPa and above, and the plasticity of the martensite after tempering is also improved.
  • austenite relieves local stress concentration as a soft phase in the early and middle stages of deformation, while martensite can occur and strengthen in the later stage of deformation. Therefore, the presence of austenite delays the initiation and propagation of cracks, and plays an important role in improving tensile strength and elongation after fracture.
  • austenite hinders crack propagation and improves crack propagation work, thereby improving impact toughness. Since the austenite in the present invention has sufficient stability, it can still exert its beneficial effect on impact toughness at -60°C.
  • the beneficial effect of the austenite structure in the present invention is closely related to its volume fraction and element enrichment degree.
  • the process parameters of the manufacturing method especially the selection of the process parameters of tempering heat treatment, most directly determine the properties of the austenite structure.
  • the mechanical properties of the core of the steel plate of the present invention are: yield strength is not less than 690MPa, tensile strength is not less than 770MPa, elongation after break is not less than 14%, impact absorption of V-shaped sample at -60°C Charpy pendulum impact test The energy is not less than 80J.
  • the present invention achieves excellent mechanical properties of the core of the steel plate, the mechanical properties at other locations in the thickness of the steel plate also reach the mechanical properties of the core. Since the present invention effectively controls the structure and properties of each position of the extra-thick steel plate, the steel plate has a high section shrinkage rate in the thickness direction, the section shrinkage rate based on the thickness direction stretching is not less than 50%, and its resistance to lamellar tearing Very good cracking performance.
  • Example 1 A 690 MPa extra-thick steel plate with excellent mechanical properties of the core, a thickness of 80 mm, and a chemical composition (content is expressed as a mass percentage) including: 0.06% C, 5.7% Mn, 0.22% Si, 0.35% Mo, 0.2 %Ni, 0.31%Cr, 0.02%Ti, S ⁇ 0.005%, P ⁇ 0.010%, Fe and other unavoidable impurity elements as residues.
  • the manufacturing method of the above steel plate is as follows:
  • converter smelting is carried out to reduce the content of S and P in molten steel to S ⁇ 0.005%, P ⁇ 0.010%; LF refining completes alloying of required mass fractions of C, Mn, Si, Mo, Ni, Ti and other elements, Afterwards, RH treatment was carried out with a vacuum of 3 mbar and a treatment time of 23 min to reduce the content of gas impurity elements in molten steel; continuous casting was used to obtain a slab with a thickness of 320 mm.
  • the slab is heated to a temperature of 1140°C and the soaking time is 60min; the heated slab is rolled, the open rolling temperature is 1005°C, the final rolling temperature is 952°C, and the rolling mill reduction procedure is 320mm-280mm-240mm-200mm-165mm- 135mm-110mm-90mm-80mm; the steel plate after rolling is water-cooled immediately, the redness temperature of the steel plate surface after cooling is 350°C, and the average cooling rate is 3.1°C/s. Quench and temper the steel plate.
  • the resulting steel sheet structure contains martensite and austenite, and the volume fraction of austenite is 6.5%.
  • Figure 1 shows the transmission electron microscope photomicrograph of the core structure of the steel plate. Martensite and austenite can be observed at intervals in the photo. The lath structure with light contrast is martensite and dark lining. The thin film-like structure is austenite.
  • the yield strength of the steel plate core is 758 MPa, the tensile strength is 842 MPa, and the elongation after break is 16%.
  • the V-shaped sample has a shock absorption energy of 135 J in the Charpy pendulum impact test at -60°C.
  • the shrinkage rate of the steel plate in the thickness direction is 63%.
  • Example 2 A 690 MPa extra-thick steel plate with excellent mechanical properties at the core, a thickness of 80 mm, and a chemical composition (content is expressed as a mass percentage) including: 0.04% C, 5.2% Mn, 0.4% Si, 0.1% Mo, 0.6 %Ni, 0.6%Cr, 0.01%Ti, S ⁇ 0.005%, P ⁇ 0.010%, Fe as the remainder and other inevitable impurity elements.
  • the manufacturing method of the above steel plate is as follows:
  • converter smelting is carried out to reduce the content of S and P in molten steel to S ⁇ 0.005%, P ⁇ 0.010%; LF refining completes alloying of required mass fractions of C, Mn, Si, Mo, Ni, Ti and other elements, Afterwards, RH treatment was carried out with a vacuum of 3 mbar and a treatment time of 20 min to reduce the content of gas impurity elements in molten steel; continuous casting was used to obtain a slab with a thickness of 320 mm.
  • the slab is heated to a temperature of 1105°C and the soaking time is 40min; the heated slab is rolled, the open rolling temperature is 1001°C, the final rolling temperature is 930°C, and the rolling mill reduction procedure is 320mm-280mm-240mm-200mm-165mm- 135mm-110mm-90mm-80mm; the steel plate after rolling is immediately water-cooled, the steel plate surface after cooling has a redness temperature of 271°C, and the average cooling rate is 4.7°C/s. Quench and temper the steel plate.
  • the resulting steel plate structure contains martensite and austenite, and the volume fraction of austenite is 10%.
  • the yield strength of the steel plate core is 741 MPa
  • the tensile strength is 821 MPa
  • the elongation after break is 17.5%
  • the V-shaped sample has a shock absorption energy of 165 J in the Charpy pendulum impact test at -60°C.
  • the steel sheet has a cross-sectional shrinkage of 71% in the thickness direction.
  • Example 3 A 690 MPa extra-thick steel plate with excellent mechanical properties of the core, a thickness of 150 mm, and the chemical composition (content is expressed as a mass percentage) include: 0.08% C, 6.0% Mn, 0.1% Si, 0.5% Mo, 0.5 %Ni, 0.2%Cr, 0.05%Ti, S ⁇ 0.005%, P ⁇ 0.010%, Fe and other inevitable impurity elements as residues.
  • the manufacturing method of the above steel plate is as follows:
  • converter smelting is carried out to reduce the content of S and P in molten steel to S ⁇ 0.005%, P ⁇ 0.010%; LF refining completes alloying of required mass fractions of C, Mn, Si, Mo, Ni, Ti and other elements, After that, RH treatment, vacuum degree 3mbar, treatment time 26min, reduce the content of gas impurity elements in molten steel; slab with a thickness of 610mm was obtained by forging after die casting.
  • the slab is heated to a temperature of 1060°C and the soaking time is 90min; the heated slab is rolled, the open rolling temperature is 1015°C, the final rolling temperature is 942°C, and the rolling mill reduction procedure is 610mm-540mm-470mm-400mm-340mm- 290mm-245mm-215mm-190mm-170mm-150mm; the rolled steel plate is immediately water-cooled.
  • the cooled steel plate surface has a redness temperature of 327°C and an average cooling rate of 1.5°C/s. Quench and temper the steel plate.
  • the resulting steel plate structure contains martensite and austenite, and the volume fraction of austenite is 4%.
  • the yield strength at the core of the steel plate is 745 MPa, the tensile strength is 819 MPa, the elongation after break is 15%, and the V-shaped sample has a shock absorption energy of 106 J in the Charpy pendulum impact test at -60°C.
  • the shrinkage rate of the steel plate in the thickness direction is 57%.
  • Example 4 Design four sets of parallel tests.
  • the component content and preparation method are basically the same as those in Example 1.
  • the difference is the open rolling temperature, as shown in Table 3 below.
  • Groups 1-2 are the open rolling temperatures within the scope of the present invention
  • Groups 3-4 are the open rolling temperatures outside the scope of the present invention.
  • the post-break elongation, low-temperature impact energy and plate thickness of the prepared steel sheet The cross-sectional shrinkage performance in the direction is poor.
  • Example 5 Design three sets of parallel tests.
  • the component content and preparation method are basically the same as in Example 2. The difference is that the cooling rate of water cooling after quenching is shown in Table 2 below.
  • the first group is the average cooling rate of water cooling after quenching within the scope of the present invention
  • the group 2-3 is the average cooling rate outside the scope of the present invention.
  • the steel plate of group 2 has poor yield strength and low-temperature impact energy performance; the steel plate of group 3 has poor low-temperature impact energy performance, and cracks caused by thermal stress appear on the steel plate.

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Abstract

一种690MPa级特厚钢板及其制造方法，化学成分的质量百分比含量为：C： 0.04-0.08％、Mn：5.2-6.0％、Si：0.1-0.4％、Mo：0.1-0.5％、Ni：0.2-0.6％、Cr：0.2-0.6％、Ti：0.01-0.05％、S：≤0.005％、P：≤0.010％以及余量的Fe和杂质。该钢板将锰作为主要合金元素，降低特厚钢板制造时镍等昂贵元素的添加量，从而降低了合金成本。采用特定的制造工艺，使制得的特厚钢板具有高强度、高塑性、高韧性的优异心部力学性能及抗层状撕裂性能，能够满足海洋工程等严苛服役环境对高性能特厚钢板的需求。

Description

一种690MPa级特厚钢板及其制造方法 技术领域
本发明涉及一种特厚钢板及其制造方法，特别涉及一种690MPa级特厚钢板及其制造方法。
背景技术
随着国家海洋发展战略的实施以及油气资源开采逐渐由陆地向深海和极地方向发展，对海洋平台的性能和结构安全性的要求越来越高。海洋平台制造所需的钢材向着高强度和高韧性方向发展，屈服强度690MPa级的高强韧海洋平台用厚板的需求量越来越大。传统的海洋工程用690MPa级特厚板的心部力学性能难以提高。为了改善钢板整体的性能均匀，通常加入大量的Ni、Mo、Cr、Cu等元素，这些元素的总添加量甚至超过4％，因此合金成本很高。此外，在制造方法中通常还需要多次淬火等工艺，制造难度大。近年来为了满足海洋工程建设对材料需求的提高，高强度级别特厚钢板的开发受到广泛关注。
专利号为201510125485.1的中国发明专利，公开了一种具有优异低温冲击韧性的低屈强比高强韧厚钢板及其制造方法，该低屈强比高强韧厚钢板化学成分中含有3.6-5.5％的Ni，成本高昂。
申请号为201610026446.0的中国专利申请，公开了一种海洋工程用高强钢板及其生产方法，采用了Nb、V微合金化方法，但由于Mn等淬透性提高元素的含量不高，钢板最大厚度未能超过100mm。
从材料开发现状来看，目前高性能的海洋工程用特厚钢板的性能还有待提升。
发明内容
发明目的：为了克服现有技术的缺陷，本发明提供一种690MPa级特厚钢板，该钢板具有优良的心部力学性能，能够满足海洋工程等严苛服役环境对高性能特厚钢板的需求。
本发明的另一目的是提供一种上述690MPa级特厚钢板的制造方法。
技术方案：本发明所述的一种690MPa级特厚钢板，化学成分的质量百分比含量为：C：0.04-0.08％、Mn：5.2-6.0％、Si：0.1-0.4％、Mo：0.1-0.5％、Ni：0.2-0.6％、Cr：0.2-0.6％、Ti：0.01-0.05％、S：≤0.005％、P：≤0.010％以及余 量的Fe和杂质。
进一步的，该钢板的厚度为80-150mm。
进一步的，该钢板的微观组织具有马氏体和奥氏体；其中，奥氏体的体积分数为4-10％。
进一步的，该钢板的心部力学性能为：屈服强度不小于690MPa，抗拉强度不小于770MPa，断后伸长率不小于14％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量不小于80J。
进一步的，该钢板基于板厚方向拉伸的断面收缩率不小于50％。
而本发明所述的一种690MPa级特厚钢板的制造方法，采用的技术方案包括如下步骤：
(1)铁水脱硫处理后进行转炉冶炼，降低钢水中S、P含量至S≤0.005％，P≤0.010％；
(2)LF精炼完成C、Mn、Si、Mo、Ni、Ti元素所需质量分数的合金化，之后进行RH处理，真空度≤4mbar，处理时间≥20min；
(3)铸造得到板坯，板坯厚度与钢板厚度的比值≥4；
(4)板坯加热，温度1060-1140℃，均热时间40-90min；
(5)对加热后的板坯进行轧制，开轧温度≤1020℃，终轧温度≥930℃，道次形变量≥10％；
(6)对轧制后的钢板立即进行水冷，冷却后的钢板表面返红温度≤360℃，平均冷却速率1-5℃/s；
(7)淬火热处理，将钢板重新加热至780-830℃，均热时间5-15min，水冷至钢板表面返红温度≤110℃温度，平均冷却速率2-8℃/s；
(8)回火热处理，将淬火后的钢板加热至610-640℃，均热时间40-70min，回火后钢板空冷至室温。
有益效果：本发明的钢板将锰作为主要合金元素，降低特厚钢板制造时镍等昂贵元素的添加量，从而降低了合金成本。采用特定的制造工艺，使制得的特厚钢板的具有高强度、高塑性、高韧性的优异心部力学性能及抗层状撕裂性能，能够满足海洋工程等严苛服役环境对高性能特厚钢板的需求。
附图说明
图1为实施例1中钢板心部组织的透射电镜显微照片。
具体实施方式
以下结合实施例对本发明的化学成分、制造工艺、组织及性能等方面进行说明。
在本发明的690MPa级特厚钢板的化学成分设计中，C能够通过间隙固溶强化显著提高组织强度，是重要的强化元素，同时也是重要的奥氏体稳定化元素，但为了保证低温冲击韧性及焊接性，需要控制其添加量在较低水平。Mn能够通过置换固溶强化提高组织强度，同时也能够显著提高奥氏体稳定性。C和Mn的适量添加能够显著提高淬透性，降低过冷奥氏体的相变温度，并获得高强度的马氏体组织。另一方面，在马氏体回火过程中，C和Mn的加入降低了形成一定量的逆转变奥氏体所需要的温度。C和Mn在逆转变奥氏体中富集，使其在低温下仍然能够保持结构稳定，成为提高本发明塑性和韧性的重要组织。需要指出的是，当本发明中回火温度低于600℃时，特别是在550℃温度附近回火时，容易造成Mn、P等元素的晶界偏聚并降低韧性。发明人充分考虑C和Mn元素在本发明中的作用机制，确定了C为0.04-0.08％、Mn为5.2-6.0％的“低碳中锰”成分设计。
Si在炼钢过程中为脱氧元素，适量的Si能够抑制Mn和P的偏聚并改善韧性。Si还能够产生固溶强化，但含量超过0.3％时会明显降低韧性。本发明将Si控制在0.1-0.4％。
Mo能够提高马氏体的回火后的强度，在一定含量范围内还能够减弱Mn的晶界偏聚从而改善韧性。本发明将Mo含量控制在0.1-0.5％，在发挥Mo作用的同时不显著增加成本。
Ni能够稳定奥氏体相、提高淬透性、降低韧脆转变温度，是提高低温韧性的有效元素，此外还有利于提高焊接性。但Ni价格昂贵，本发明将Ni含量控制在0.2-0.6％，在不显著增加成本的同时充分发挥Ni元素的有益作用。
Cr能够产生明显的固溶强化作用，有利于提高强度，能够改善耐蚀性。但在本发明添加较多Mn元素的情况下，Cr含量过高容易在回火时在晶界处形成Cr、Mn的碳化物，降低晶界对裂纹扩展的阻碍，导致塑性和韧性降低。本发明将Cr的含量范围控制在0.2-0.6％的适宜范围。
本发明中加入微量的Ti，能够通过细小而弥散的第二相析出形式阻碍高温 下的晶界迁移，从而细化晶粒并改善力学性能，加入量控制在0.01-0.05％的范围内。
严格控制P与S的含量，在本发明已添加中等含量Mn元素的情况下，S易与Mn形成MnS并降低塑性。P容易在晶界偏聚，降低晶界抗裂纹扩展能力，从而降低韧性。本发明要求S≤0.005％，P≤0.010％。
本发明的剩余物为Fe，然而，在通常的制造过程中会不可避免地从原料或周围环境中引入杂质。由于这些杂质对于本领域技术人员而言是显而易见的，所以其名称及含量不具体地记载于本说明书中。
在本发明制造方法中，通过铁水脱硫处理后进行转炉冶炼降低S和P的含量至S≤0.005％，P≤0.010％，并通过足够高真空度(真空度≤4mbar)、足够长真空时间(处理时间≥20min)的RH处理以降低气体杂质元素的含量，而C、Mn、Si、Mo、Ni、Ti等合金的加入则通过LF精炼完成，因此，能够实现高纯净度的冶炼效果。
板坯铸造可以采用连铸或模铸+锻造的方式，获得不同尺寸规格的板坯。板坯厚度≤320mm的情况可以通过连铸方式获得，生产效率高；更大厚度的板坯(＞320mm)可以通过模铸+锻造的方式。为了实现本发明所需要的心部力学性能，需要足够的轧制总形变量作为必要条件。本发明中要求板坯厚度与钢板厚度的比值≥4，能够保证轧制总行变量≥75％。钢板厚度为80mm时，所需的板坯厚度≥320mm；钢板厚度为150mm时，所需的板坯厚度≥600mm。所获得的板坯通过轧制及热处理工艺获得所需要的组织及性能。
在本发明成分范围内，钢的Ac3温度不高于770℃。板坯加热到1060-1140℃时形成高温奥氏体组织，同时C、Mn等合金元素通过扩散方式均匀化。在板坯心部温度接近表面温度并继续保温的均热过程中，整个板坯实现奥氏体的均匀化，均温时间40-90min能够保证元素扩散均匀。在1140℃及以下温度，Ti的第二相颗粒能够起到阻碍晶粒长大的作用。但温度低于1060℃时元素扩散过慢，奥氏体均匀化效率过低。
本发明对加热后的板坯在930℃及以上进行再结晶轧制以细化晶粒。开轧温度≤1020℃能够避免再结晶后的晶粒长大速率过快。道次形变量≥10％能够使奥氏体形变后具有足够的畸变能累积，以保证再结晶细化效果。
钢板轧制后，为了避免轧制形变后细化的再结晶晶粒过度长大，需要对轧制后的钢板立即进行水冷。水冷过程中还将发生马氏体转变。由于本发明添加了足量的Mn元素，马氏体转变临界冷却速率低于1℃/s，冷却速率较低的情况下也能够得到马氏体组织。在钢板厚度较大的情况下，心部冷却通常明显慢于表面冷却，本发明的成分设计能够保证在厚度为80-150mm的钢板的心部也发生马氏体转变。但冷速过高容易导致过高的钢板热应力，甚至导致钢板开裂，因此本发明控制平均冷却速率不高于5℃/s。轧制后冷却的钢板表面返红温度选择在360℃及以下，能够避免冷却过程中发生明显的元素偏聚，并抑制粗大碳化物的析出，同时这一温度也低于本发明成分下的马氏体转变开始温度。本发明选择的轧制后冷却工艺能够为后续的热处理工艺提供合适的前驱组织。
本发明对钢板进行淬火+回火热处理。780-830℃的淬火温度高于Ac3，均热得到奥氏体组织。由于轧制后对钢板进行了水冷，避免了冷却过程中元素偏聚及粗大碳化物的形成，因此奥氏体内元素均匀化时间大大缩短。本发明选择淬火加热的均热时间为5-15min，在保证奥氏体均匀化的同时还能够有效细化晶粒尺寸，有利于钢板力学性能的提高。淬火冷却速率的选择与轧制后冷却速率的选择理由相同，控制在2-8℃/s。但是，淬火的钢板表面返红温度要求在110℃及以下，这一温度低于本发明成分下的马氏体转变终止温度，能够确保钢板整体得到高强度的淬火马氏体组织。
对淬火后的钢板进行回火热处理。在回火温度610-640℃、均热时间40-70min的回火过程中，除了改善马氏体的强韧性匹配外，还能够得到体积分数为4-10％的逆转变奥氏体，主要呈薄膜状，分布于马氏体板条之间。回火过程中，C和Mn等奥氏体稳定化元素在奥氏体内富集，提高了奥氏体的稳定性，在回火后空冷至室温的过程中，甚至在更低温度的情况下，奥氏体仍然能够保持晶体结构的稳定，而不发生马氏体转变。回火后钢板空冷，还能够降低特厚钢板的热应力，改善钢板质量。
本发明热处理后，在钢板整个厚度方向，特别是在钢板心部得到了马氏体+奥氏体的组织。在拉伸形变过程中，马氏体基体提供了690MPa及以上的屈服强度，回火后的马氏体的塑性也得到改善。奥氏体在形变的前期和中期作为一种软相缓解局部应力集中，而在形变的后期能够发生马氏体并产生强化作用。因此， 奥氏体的存在延缓了裂纹的萌生和扩展，起到了提高抗拉强度和断后伸长率的重要作用。在冲击形变过程中，奥氏体的存在阻碍了裂纹扩展，提高了裂纹扩展功，从而提高了冲击韧性。由于本发明中的奥氏体具有足够的稳定性，其在-60℃下仍然能够发挥对冲击韧性的有益作用。本发明中奥氏体组织的有益作用，与其体积分数和元素富集程度密切相关，制造方法的工艺参数，特别是回火热处理的工艺参数选择最为直接地决定了奥氏体组织的性质。
具体的，本发明钢板的心部力学性能为：屈服强度不小于690MPa，抗拉强度不小于770MPa，断后伸长率不小于14％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量不小于80J。
本发明中力学性能指标的定义依照标准GB/T228.1、GB/T229和GB/T5313，由于这些技术指标的定义对于本领域技术人员而言是明确的，所以在本说明书中不作过多说明。
需要指出的是，本发明在实现了优异的钢板心部力学性能的同时，钢板厚度其它位置的力学性能同样达到心部的力学性能。由于本发明对特厚钢板各个位置的组织和性能进行了有效控制，钢板具有很高的厚度方向断面收缩率，其基于板厚方向拉伸的断面收缩率不小于50％，其抗层状撕裂性能十分优异。
下面，以具体实施例上述钢板及其制造方法做进一步说明。
实施例1：一种心部力学性能优异的690MPa级特厚钢板，厚度为80mm，化学成分(含量表示为质量百分数)包括：0.06％C，5.7％Mn，0.22％Si，0.35％Mo，0.2％Ni，0.31％Cr，0.02％Ti，S≤0.005％，P≤0.010％，作为剩余物的Fe以及其他不可避免的杂质元素。
上述钢板的制造方法如下：
铁水脱硫处理后进行转炉冶炼，降低钢水中S、P含量至S≤0.005％，P≤0.010％；LF精炼完成C、Mn、Si、Mo、Ni、Ti等元素所需质量分数的合金化，之后进行RH处理，真空度3mbar，处理时间23min，降低钢水中的气体杂质元素含量；采用连铸方式得到厚度为320mm的板坯。板坯加热到温度1140℃，均热时间60min；对加热后的板坯进行轧制，开轧温度1005℃，终轧温度952℃，轧机压下规程为320mm-280mm-240mm-200mm-165mm-135mm-110mm-90mm-80mm；对轧制后的钢板立即进行水冷，冷却后的钢板表面返红温度350℃，平均冷却速 率3.1℃/s。对钢板进行淬火+回火热处理。淬火温度810℃，均热时间10min，水冷至钢板表面返红温度77℃温度，平均冷却速率2.9℃/s；回火温度626℃，均热时间55min，回火后钢板空冷至室温。
得到的钢板组织含有马氏体和奥氏体，奥氏体的体积分数为6.5％。图1所示为钢板心部组织的透射电镜显微照片，照片中可观察到间隔分布的马氏体和奥氏体，其中浅色衬度的板条状组织为马氏体，深色衬度的薄膜状组织为奥氏体。钢板心部的屈服强度758MPa，抗拉强度842MPa，断后伸长率16％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量135J。钢板的厚度方向断面收缩率63％。
实施例2：一种心部力学性能优异的690MPa级特厚钢板，厚度为80mm，化学成分(含量表示为质量百分数)包括：0.04％C，5.2％Mn，0.4％Si，0.1％Mo，0.6％Ni，0.6％Cr，0.01％Ti，S≤0.005％，P≤0.010％，作为剩余物的Fe以及其他不可避免的杂质元素。
上述钢板的制造方法如下：
铁水脱硫处理后进行转炉冶炼，降低钢水中S、P含量至S≤0.005％，P≤0.010％；LF精炼完成C、Mn、Si、Mo、Ni、Ti等元素所需质量分数的合金化，之后进行RH处理，真空度3mbar，处理时间20min，降低钢水中的气体杂质元素含量；采用连铸方式得到厚度为320mm的板坯。板坯加热到温度1105℃，均热时间40min；对加热后的板坯进行轧制，开轧温度1001℃，终轧温度930℃，轧机压下规程为320mm-280mm-240mm-200mm-165mm-135mm-110mm-90mm-80mm；对轧制后的钢板立即进行水冷，冷却后的钢板表面返红温度271℃，平均冷却速率4.7℃/s。对钢板进行淬火+回火热处理。淬火温度830℃，均热时间5min，水冷至钢板表面返红温度51℃温度，平均冷却速率4.2℃/s；回火温度640℃，均热时间40min，回火后钢板空冷至室温。
得到的钢板组织含有马氏体和奥氏体，奥氏体的体积分数为10％。钢板心部的屈服强度741MPa，抗拉强度821MPa，断后伸长率17.5％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量165J。钢板的厚度方向断面收缩率71％。
实施例3：一种心部力学性能优异的690MPa级特厚钢板，厚度为150mm，化学成分(含量表示为质量百分数)包括：0.08％C，6.0％Mn，0.1％Si，0.5％Mo，0.5％Ni，0.2％Cr，0.05％Ti，S≤0.005％，P≤0.010％，作为剩余物的Fe以及其他不可避免 的杂质元素。
上述钢板的制造方法如下：
铁水脱硫处理后进行转炉冶炼，降低钢水中S、P含量至S≤0.005％，P≤0.010％；LF精炼完成C、Mn、Si、Mo、Ni、Ti等元素所需质量分数的合金化，之后进行RH处理，真空度3mbar，处理时间26min，降低钢水中的气体杂质元素含量；采用模铸后锻造方式得到厚度为610mm的板坯。板坯加热到温度1060℃，均热时间90min；对加热后的板坯进行轧制，开轧温度1015℃，终轧温度942℃，轧机压下规程为610mm-540mm-470mm-400mm-340mm-290mm-245mm-215mm-190mm-170mm-150mm；对轧制后的钢板立即进行水冷，冷却后的钢板表面返红温度327℃，平均冷却速率1.5℃/s。对钢板进行淬火+回火热处理。淬火温度780℃，均热时间15min，水冷至钢板表面返红温度102℃温度，平均冷却速率1.2℃/s；回火温度610℃，均热时间70min，回火后钢板空冷至室温。
得到的钢板组织含有马氏体和奥氏体，奥氏体的体积分数为4％。钢板心部的屈服强度745MPa，抗拉强度819MPa，断后伸长率15％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量106J。钢板的厚度方向断面收缩率57％。
实施例4：设计4组平行试验，组分含量及制备方法与实施例1基本相同，不同之处在于开轧温度，具体如下表3所示。
表1实施例4钢板的力学性能
Figure PCTCN2019122904-appb-000001
通过表1可知，第1-2组为本发明范围内的开轧温度，而3-4组为本发明范围外的开轧温度，其制备钢板的断后伸长率、低温冲击功以及板厚方向的断面收缩率性能差。
实施例5：设计3组平行试验，组分含量及制备方法与实施例2基本相同，不同之处在于，淬火后水冷的冷却速度，具体如下表2所示。
表2实施例5钢板的力学性能
Figure PCTCN2019122904-appb-000002
通过表2可知，第1组为本发明范围内的淬火后水冷的平均冷却速度，而2-3组为本发明范围外的平均冷却速度。组号2的钢板屈服强度和低温冲击功性能差；组号3的钢板低温冲击功性能差，而且钢板上出现了热应力导致的裂纹。

Claims (10)

  1. 一种690MPa级特厚钢板，其特征在于，化学成分的质量百分比含量为：C：0.04-0.08％、Mn：5.2-6.0％、Si：0.1-0.4％、Mo：0.1-0.5％、Ni：0.2-0.6％、Cr：0.2-0.6％、Ti：0.01-0.05％、S：≤0.005％、P：≤0.010％以及余量的Fe和杂质。
  2. 根据权利要求1所述的690MPa级特厚钢板，其特征在于，该钢板的厚度为80-150mm。
  3. 根据权利要求1所述的690MPa级特厚钢板，其特征在于，该钢板的微观组织具有马氏体和奥氏体；其中，奥氏体的体积分数为4-10％。
  4. 根据权利要求3所述的690MPa级特厚钢板，其特征在于，所述奥氏体呈薄膜状，奥氏体分布于马氏体板条之间。
  5. 根据权利要求1所述的690MPa级特厚钢板，其特征在于，该钢板的心部力学性能为：屈服强度不小于690MPa，抗拉强度不小于770MPa，断后伸长率不小于14％，V型试样-60℃夏比摆锤冲击试验冲击吸收能量不小于80J。
  6. 根据权利要求1所述的690MPa级特厚钢板，其特征在于，该钢板基于板厚方向拉伸的断面收缩率不小于50％。
  7. 一种根据权利要求1-6任一项所述的690MPa级特厚钢板的制备方法，其特征在于，包括如下步骤：
    (1)铁水脱硫处理后进行转炉冶炼，降低钢水中S、P含量至S≤0.005％，P≤0.010％；
    (2)LF精炼完成C、Mn、Si、Mo、Ni、Ti元素所需质量分数的合金化，之后进行RH处理，真空度≤4mbar，处理时间≥20min；
    (3)铸造得到板坯，板坯厚度与钢板厚度的比值≥4；
    (4)板坯加热，温度1060-1140℃；
    (5)对加热后的板坯进行轧制，开轧温度≤1020℃，终轧温度≥930℃，道次形变量≥10％；
    (6)对轧制后的钢板立即进行水冷，冷却后的钢板表面返红温度≤360℃，平均冷却速率1-5℃/s；
    (7)淬火热处理，将钢板重新加热至780-830℃，均热时间5-15min，水冷至钢板表面返红温度≤110℃温度，平均冷却速率2-8℃/s；
    (8)回火热处理，将淬火后的钢板加热至610-640℃，均热时间40-70min，回火后钢板空冷至室温。
  8. 根据权利要求6所述的制备方法，其特征在于，所述步骤(3)中板坯采用连铸板坯或者模铸后锻造板坯。
  9. 根据权利要求7所述的制备方法，其特征在于，当板坯厚度≤320mm时，采用连铸板坯；当板坯厚度＞320mm时，采用模铸后锻造板坯。
  10. 根据权利要求6所述的制备方法，其特征在于，所述步骤(4)中，板坯的均热时间为40-90min。
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Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN112705582A (zh) * 2021-01-28 2021-04-27 李艳峰 一种铅箔成形方法

Families Citing this family (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN109652733B (zh) * 2019-01-07 2021-01-26 南京钢铁股份有限公司 一种690MPa级特厚钢板及其制造方法
CN111778450A (zh) * 2020-06-24 2020-10-16 南京钢铁股份有限公司 一种800MPa工程机械用中锰中厚钢及其制造方法
CN112359289B (zh) * 2020-11-23 2021-12-21 马鞍山钢铁股份有限公司 一种超厚q355级良好低温韧性热轧h型钢及生产方法
CN112899584A (zh) * 2021-01-15 2021-06-04 南京钢铁股份有限公司 超低温l型钢及其制造方法
CN115323251B (zh) * 2022-08-24 2023-06-27 东北大学 一种超厚高强韧高均质水电用特厚钢板及其制造方法

Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN102226255A (zh) * 2011-06-08 2011-10-26 江苏省沙钢钢铁研究院有限公司 屈服强度690MPa高强韧钢板及其制备工艺
CN103031498A (zh) * 2012-12-17 2013-04-10 南京钢铁股份有限公司 低压缩比特厚超高强应变时效的海洋工程钢板的制造方法
CN104711489A (zh) * 2015-03-03 2015-06-17 山东钢铁股份有限公司 一种桩腿半弦管用高韧性钢及其制造方法
US20160312327A1 (en) * 2013-12-12 2016-10-27 Jfe Steel Corporation Steel plate and method for manufacturing same (as amended)
CN106319380A (zh) * 2015-06-16 2017-01-11 鞍钢股份有限公司 一种低压缩比690MPa级特厚钢板及其生产方法
CN107312981A (zh) * 2017-06-13 2017-11-03 南京钢铁股份有限公司 一种低屈强比高强韧厚规格钢板及其制造方法
CN109652733A (zh) * 2019-01-07 2019-04-19 南京钢铁股份有限公司 一种690MPa级特厚钢板及其制造方法

Family Cites Families (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP3228986B2 (ja) * 1992-02-12 2001-11-12 新日本製鐵株式会社 高張力鋼板の製造方法
JP4283757B2 (ja) 2004-11-05 2009-06-24 株式会社神戸製鋼所 厚鋼板およびその製造方法
JP6237343B2 (ja) 2014-02-28 2017-11-29 新日鐵住金株式会社 高清浄鋼の溶製方法

Patent Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN102226255A (zh) * 2011-06-08 2011-10-26 江苏省沙钢钢铁研究院有限公司 屈服强度690MPa高强韧钢板及其制备工艺
CN103031498A (zh) * 2012-12-17 2013-04-10 南京钢铁股份有限公司 低压缩比特厚超高强应变时效的海洋工程钢板的制造方法
US20160312327A1 (en) * 2013-12-12 2016-10-27 Jfe Steel Corporation Steel plate and method for manufacturing same (as amended)
CN104711489A (zh) * 2015-03-03 2015-06-17 山东钢铁股份有限公司 一种桩腿半弦管用高韧性钢及其制造方法
CN106319380A (zh) * 2015-06-16 2017-01-11 鞍钢股份有限公司 一种低压缩比690MPa级特厚钢板及其生产方法
CN107312981A (zh) * 2017-06-13 2017-11-03 南京钢铁股份有限公司 一种低屈强比高强韧厚规格钢板及其制造方法
CN109652733A (zh) * 2019-01-07 2019-04-19 南京钢铁股份有限公司 一种690MPa级特厚钢板及其制造方法

Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN112705582A (zh) * 2021-01-28 2021-04-27 李艳峰 一种铅箔成形方法

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