WO2025002365A1 - 一种管线钢及其制造方法 - Google Patents

一种管线钢及其制造方法 Download PDF

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Publication number
WO2025002365A1
WO2025002365A1 PCT/CN2024/102421 CN2024102421W WO2025002365A1 WO 2025002365 A1 WO2025002365 A1 WO 2025002365A1 CN 2024102421 W CN2024102421 W CN 2024102421W WO 2025002365 A1 WO2025002365 A1 WO 2025002365A1
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steel
content
pipeline steel
pipeline
resistance
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English (en)
French (fr)
Inventor
张豪臻
章传国
王波
孙磊磊
李龙
梅峰
陈国锋
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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 DE112024002762.8T priority Critical patent/DE112024002762T5/de
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    • 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
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/74—Temperature control, e.g. by cooling or heating the rolls or the product
    • B21B37/76—Cooling control on the run-out table
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C21C7/0075—Treating in a ladle furnace, e.g. up-/reheating of molten steel within the ladle
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C21C7/04—Removing impurities by adding a treating agent
    • C21C7/064—Dephosphorising; Desulfurising
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C21C7/10—Handling in a vacuum
    • 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
    • C21D11/00—Process control or regulation for heat treatments
    • 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
    • C21D11/00—Process control or regulation for heat treatments
    • C21D11/005—Process control or regulation for heat treatments for cooling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—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
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
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    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
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    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • 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
    • C22C33/06—Making ferrous alloys by melting using master alloys
    • 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
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • 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
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—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
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • 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
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the invention belongs to the field of metal materials and their manufacturing, and in particular relates to pipeline steel and a manufacturing method thereof.
  • H2S -containing acidic oil and gas fields With the continuous development of conventional oil and gas fields, the reserves are constantly decreasing, and the development ratio of H2S -containing acidic oil and gas fields is further increasing, which requires the pipeline to have acid resistance.
  • the designers of different acid-resistant pipeline projects will formulate personalized acid resistance-related indicators according to the service conditions of the pipeline, including NACE standard resistance to hydrogen-induced cracking (HIC) performance, resistance to hydrogen sulfide stress corrosion cracking (SSC) performance, and strict low-power rating and flaw detection requirements.
  • HIC hydrogen-induced cracking
  • SSC hydrogen sulfide stress corrosion cracking
  • Chinese patent CN202010552764.7 discloses "a magnesium-containing X65 pipeline steel with excellent acid resistance and production method".
  • the components and weight percentage of the pipeline steel are: C: 0.03-0.05%, Si: 0.17-0.20%, Mn: 0.40-0.70%, Ni: 0.12-0.15%, Cr: 0.45-0.48%, Mo: 0.10-0.15%, Mg: 0.002-0.004%, Nb: 0.06-0.09%, Ti: 0.01-0.02%, P ⁇ 0.008%, S ⁇ 0.010%, and the rest are Fe and unavoidable impurities.
  • the patent adopts the Mg treatment method for inclusion treatment, and the Mn content in the chemical composition is relatively low.
  • the acid-resistant pipeline steel with X65 strength level is obtained.
  • the use of Mg to treat inclusions makes it difficult to control the stable addition, and the control effect on inclusions such as MnS is poor, which has an adverse effect on acid resistance and toughness.
  • the obtained acid-resistant pipeline steel has a Charpy impact energy of ⁇ 200J at -20°C and a DWTT SA% of ⁇ 90% at -15°C.
  • Ci 0.050-0.070%
  • Si 0.15-0.25%
  • Mn 1.35-1.45%
  • Nb 0.025-0.045%
  • Ti 0.008-0.025%
  • Mo 0.06-0 .10%
  • Cr 0.08 ⁇ 0.18%
  • Als 0.015 ⁇ 0.045%
  • P ⁇ 0.015% S ⁇ 0.0015%
  • N ⁇ 0.006% Due to the increase of C content, the segregation degree in the steel will be significantly increased, the HIC sensitivity of the steel will be increased, and the low-temperature toughness of the steel will be reduced to a certain extent.
  • the Charpy impact energy of the obtained steel at -40°C is ⁇ 300J
  • the DWTT SA% at -30°C is ⁇ 85%.
  • Chinese patent CN202011197059.6 discloses "A X60/X65 grade acid-resistant pipeline steel and its preparation method", whose chemical composition and mass percentage are as follows: C: 0.02-0.06%, Si: 0.10-0.20%, Mn ⁇ 0.90%, Alt: 0.020-0.050%, Nb: 0.060-0.080%, Ti ⁇ 0.020%, Cr ⁇ 0.30%, Ni ⁇ 0.20%, Cu ⁇ 0.40%, P ⁇ 0.012%, S ⁇ 0.002%, N ⁇ 0.0060%, B ⁇ 0.0005%, Mn+Cr ⁇ 1.2%, the rest is iron and unavoidable
  • the patented alloy composition design adds more Nb (0.060-0.080%).
  • the high Nb content increases the tendency of Nb-containing particles in the steel to gather in the core, and the probability of HIC cracks in the core of the steel plate increases accordingly.
  • the object of the present invention is to provide a high-toughness X60/X65 grade acid-resistant pipeline steel and a manufacturing method thereof.
  • the pipeline steel can reach the X60/X65 strength level and has good low-temperature toughness and excellent acid resistance (such as HIC resistance and SSC resistance).
  • the pipeline steel has an R t0.5 yield strength of 430-570 MPa, a tensile strength of 520-690 MPa, a yield strength ratio of ⁇ 0.92, an elongation A50 of ⁇ 30%, a hardness HV10 of ⁇ 220, a -49°C Charpy impact energy of ⁇ 400 J, and a -35°C DWTT SA% of ⁇ 90%.
  • the present invention provides a pipeline steel, which, in addition to Fe and unavoidable impurities, also contains the following chemical elements by mass percentage: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0 ⁇ P ⁇ 0.008%, 0 ⁇ S ⁇ 0.0010%, Cu: 0.05-0.35%, Ni: 0.05-0.30%, Cr: 0.15-0.30%, Mo: 0.02-0.09%, N b: 0.025-0.055%, Ti: 0.005-0.020%, Ca: 0.0010-0.0040%, Alt: 0.010-0.040%, 0 ⁇ B ⁇ 0.0004%, 0 ⁇ O ⁇ 0.0026%, 0 ⁇ N ⁇ 0.0050%, 0 ⁇ H ⁇ 0.0002%, 0 ⁇ Sn ⁇ 0.0050%, 0 ⁇ Sb ⁇ 0.0025%, 0 ⁇ Bi ⁇ 0.0030%, 0 ⁇ Pb ⁇ 0.0050%, 0 ⁇ As ⁇ 0.0030%, and satisfying:
  • the pipeline steel contains the following chemical elements by mass percentage: C: 0.015-0.040%, Si: 0.10-0.30%, Mn: 1.18-1.45%, 0 ⁇ P ⁇ 0.008%, 0 ⁇ S ⁇ 0.0010%, Cu: 0.05-0.35%, Ni: 0.05 ⁇ 0.30%, Cr: 0.15 ⁇ 0.30%, Mo: 0.02 ⁇ 0.09%, Nb: 0.025 ⁇ 0.055%, Ti: 0.005 ⁇ 0.020%, Ca: 0.0010 ⁇ 0.0040%, Alt: 0.010 ⁇ 0.040%, 0 ⁇ B ⁇ 0.0004%, 0 ⁇ O ⁇ 0.0026%, 0 ⁇ N ⁇ 0.0050%, 0 ⁇ H ⁇ 0.0002%, 0 ⁇ Sn ⁇ 0.0050%, 0 ⁇ Sb ⁇ 0.0025%, 0 ⁇ Bi ⁇ 0.0030%, 0 ⁇ Pb ⁇ 0.0050%, 0 ⁇ As ⁇ 0.0030%, the balance is Fe and unavoidable impurities, and meets the following requirements:
  • C is the most economical strengthening element in steel, which improves the strength of steel through interstitial solid solution strengthening. Increasing the carbon content can greatly improve the hardenability of steel, reduce the addition of other precious alloying elements, and reduce production costs, but the increase in C content is unfavorable to the ductility, toughness, weldability and corrosion resistance of steel, especially prone to aggravate core segregation, and has a significant adverse effect on acid resistance. Therefore, the present invention adopts an ultra-low C design, and its mass percentage content is controlled at 0.015-0.040%, preferably 0.017-0.032%.
  • Si is a solid solution strengthening element and also a deoxidizing element in steel. However, too high Si content will have an adverse effect on surface quality and welding performance. If the Si content exceeds 0.30%, toughness may decrease. Therefore, in the present invention, the mass percentage content of Si is controlled to 0.10-0.30%, preferably 0.18-0.27%.
  • Mn element improves the strength of steel by solid solution strengthening. It is the most important and economical strengthening element in steel to compensate for the strength loss caused by the reduction of C content. Mn helps to obtain fine phase transformation products and also helps to control oxygen and sulfur during the steelmaking process. If the Mn content is too low, it will be difficult to achieve the target strength level, but if the Mn content is too high, it will also aggravate the center segregation. Therefore, in the present invention, the mass percentage content of Mn is controlled to 1.18-1.45%, preferably 1.25-1.43%.
  • Nb+Ti Adding too much Nb and Ti will easily lead to the aggregation of coarser Nb/Ti-containing particles in the steel plate, and easily form long strip-shaped segregation distributed along the rolling direction in the core.
  • the long strip-shaped segregation is a hard phase and its hardness is higher than that of the matrix, which increases the hydrogen-induced cracking sensitivity of the organization, causing hydrogen to be pinned around it and no longer diffuse. Therefore, it is easy to become a hydrogen aggregation site and also easy to become the starting point of hydrogen-induced cracking, significantly increasing the HIC cracking tendency of the core and deteriorating the acid resistance of the steel. Therefore, in the present invention, the mass percentage content of Nb and Ti elements is controlled to satisfy Nb+Ti ⁇ 0.07%, preferably 0.033-0.069%, and more preferably 0.033-0.065%.
  • TiN and NbC are important second phases in steel, which have an important influence on strength and toughness.
  • the second phase particles are relatively coarse and unevenly distributed, and are easy to gather at the core segregation, causing stress concentration, triggering brittle fracture, and causing a sharp decrease in the low-temperature toughness of the steel, significantly reducing the toughness of the steel and the hot zone of the weld, and also becoming a hard phase that causes hydrogen-induced cracking.
  • the mass percentage content of Nb and Ti elements is controlled to satisfy 0.7 ⁇ (Nb+27.5*Ti)/(5C+27*N) ⁇ 5.0, preferably 0.81 ⁇ (Nb+27.5*Ti)/(5C+27*N) ⁇ 3.69, and more preferably 1.13 ⁇ (Nb+27.5*Ti)/(5C+27*N) ⁇ 2.70.
  • Mo is a strong hardenability element, which can significantly delay the ferrite phase transformation, inhibit the formation of ferrite and pearlite, effectively promote the bainite transformation and play a role in strengthening the matrix, and obtain a finer structure.
  • a certain amount of Mo has obvious benefits for improving the structural properties of steel, but when the Mo content is too high, the plasticity of the steel is reduced, and the cost of Mo is high. Therefore, in the present invention, the mass percentage content of Mo is controlled at a low level of 0.02 to 0.09%.
  • the mass percentage content of Cu is controlled to 0.05-0.35%, preferably 0.15-0.33%; the mass percentage content of Ni is controlled to 0.05-0.30%, preferably 0.10-0.22%.
  • Cu+Ni+Cr+Mo The addition of Cu, Ni, Cr, and Mo elements will increase the production cost, and if the content of Cu, Ni, Cr, and Mo is too high, it will increase the hot brittleness tendency of the steel, and at the same time promote the formation of MA particles, increase the proportion of hard phases in the steel, and thus become a hydrogen trap, resulting in a series of problems such as hydrogen embrittlement and brittle phases, thereby reducing the HIC resistance and low-temperature toughness, affecting the performance of the steel.
  • the mass percentage content of the Cu, Ni, Cr, and Mo elements in the present invention is controlled to meet Cu+Ni+Cr+Mo ⁇ 0.80%, preferably 0.34 ⁇ Cu+Ni+Cr+Mo ⁇ 0.79%.
  • S, P S and P are the main impurity elements in steel.
  • P easily causes cold brittleness of steel, and S easily causes hot brittleness, which leads to unstable performance of steel.
  • the mass percentage of S and P elements in the present invention is controlled to be S ⁇ 0.0010% and P ⁇ 0.0080%.
  • Ca treatment can control the morphology of sulfides, improve the anisotropy of the steel plate, and increase low-temperature toughness; in addition, when the S content is too high, the appropriate Ca/S ratio is controlled to ensure the HIC resistance of the steel, but too high a Ca content is likely to increase the inclusions such as CaO in the steel.
  • the mass percentage of Ca is controlled to be 0.0010-0.0040%, preferably 0.0021-0.0035%, and Ca/S is controlled to be ⁇ 2.
  • the Ca/S ratio is 3-13, and more preferably 4.1-8.7.
  • Al is used for deoxidation of steel. A proper amount of Al is also beneficial for grain refinement and improving toughness. However, if the total aluminum (Alt) content is greater than 0.04%, coarse precipitates may be formed, thereby reducing the acid resistance and low-temperature toughness of the steel. Therefore, the mass percentage content of Alt is controlled within 0.010-0.040% in the present invention.
  • N, O, H For low-temperature toughness, O, N and H elements may all cause brittle fracture, especially at low temperatures. The presence of O and N will form O compounds and N compounds, which may aggregate at the grain boundaries, leading to intergranular brittleness. At the same time, H may also penetrate into the grains, causing the lattice to deform and form holes, leading to brittle fracture. Therefore, the low-temperature toughness of steel is related to the content of O, N, and H. Excessive or improper content may lead to brittle fracture. For acid resistance, the effects of O, N and H are also significant. In an acidic environment, O will react with the iron element in the steel to oxidize, causing the steel to lose a certain degree of corrosion resistance.
  • N may form N compounds with the chromium element in the steel, resulting in a decrease in the corrosion resistance of the steel.
  • H will promote stress corrosion cracking of the steel, which will also have a certain impact on the acid resistance of the steel. Therefore, in the present invention, the mass percentage content of N, O, and H elements is controlled to be 0 ⁇ O ⁇ 0.0026%, 0 ⁇ N ⁇ 0.0050%, 0 ⁇ H ⁇ 0.0002%, and 0 ⁇ N+O+H ⁇ 0.0070%, preferably 0.0024% ⁇ N+O+H ⁇ 0.0065%.
  • B The main function of B element in steel is to increase the hardenability and strength of steel, thereby saving other rare and precious metals, but the addition of B element will have a significant adverse effect on the low-temperature toughness of steel. Therefore, in the present invention, the mass percentage content of B element is controlled to 0 ⁇ B ⁇ 0.0004%.
  • Sn, Sb, Bi, Pb, As When the content of Sn, Sb, Bi, Pb, and As in steel exceeds a certain limit, the strength of the steel will be reduced, and it is easy to cause serious segregation, forming segregation at the grain boundary, causing brittle cracks to start from the grain boundary and extend along the grain boundary until complete fracture, significantly reducing the low-temperature toughness of the steel; in addition, excessive Sn content in steel will significantly reduce the steel's resistance to HIC hydrogen-induced cracking, because Sn will promote the diffusion and aggregation of hydrogen, thereby accelerating the occurrence of hydrogen-induced cracking, and Pb also has a similar effect.
  • the mass percentage content of Sn, Sb, Bi, Pb, and As elements is controlled to 0 ⁇ Sn ⁇ 0.0050%, 0 ⁇ Sb ⁇ 0.0025%, 0 ⁇ Bi ⁇ 0.0030%, 0 ⁇ Pb ⁇ 0.0050%, and 0 ⁇ As ⁇ 0.0030%.
  • the present invention adopts the above-designed component system, comprehensively considering the effects of easily segregated elements such as C and S in steel, Cu, Ni, Cr, Mo and Nb, Ti microalloying elements, and N, O, H, B and Sn, Sb, Bi, Pb, As and other elements on the acid resistance and low-temperature toughness of steel, and proposes control requirements for the hydrogen-induced cracking sensitivity coefficient Ks based on the degree of segregation of each element and its relationship with the HIC resistance, so as to improve the acid resistance of steel.
  • Nb+Ti is controlled to be ⁇ 0.07%, 0.7 ⁇ (Nb+27.5*Ti)/(5C+27*N) ⁇ 5.0.
  • Cu+Ni+Cr+Mo 0.0%
  • the inhibitory effect on massive ferrite is reduced, so that a small amount of massive ferrite is contained in the steel structure (the content of massive ferrite is 3 to 10 volume %), thereby inhibiting the occurrence of cracks, so that the pipeline steel has good low-temperature toughness on the basis of high strength and acid resistance.
  • the microstructure of the pipeline steel is granular bainite+blocky ferrite.
  • the content of the massive ferrite is 3 to 10 volume %, more preferably 4.6 to 8.5 volume %.
  • the grain size of the blocky ferrite is ⁇ 10, more preferably 10.5-11.
  • the content of the granular bainite is 90 to 97 volume %, more preferably 91.5 to 95.4 volume %.
  • the granular bainite is uniformly refined granular bainite.
  • the pipeline steel has an R t0.5 yield strength of 430 MPa or more, preferably 450-570 MPa, and more preferably 510-570 MPa; a tensile strength of 520 MPa or more, preferably 540-690 MPa, and more preferably 615-690 MPa; a yield strength ratio of ⁇ 0.92, preferably ⁇ 0.87, and more preferably 0.81-0.87; an elongation A50 of ⁇ 30%, preferably 30%-47%, and more preferably 40%-47%; a hardness HV10 of ⁇ 220, and preferably 195-220; a Charpy impact energy at -49°C of ⁇ 400 J, preferably ⁇ 430 J, and more preferably 430 J-500 J; and a DWTT SA% of -35°C of ⁇ 90%, and preferably 90-98%.
  • the pipeline steel meets at least one of the following properties: no cracks in the HIC resistance test conducted according to the NACE TM0284-2016 standard "Test method for evaluation of hydrogen-induced cracking resistance of steel for pipelines and pressure vessels"; no cracks in the SSC resistance test conducted according to the NACE TM 0177-2016 standard "Test method for resistance of metallic materials to sulfide stress corrosion cracking in pressurized sulfur environments”.
  • the present invention also provides a method for manufacturing the above pipeline steel, which comprises the following steps:
  • the molten iron in step 1) is smelted to control the C content of the tapped steel to be ⁇ 0.035% and the S content of the tapped steel to be ⁇ 0.0080%;
  • the steel produced in step 2) is subjected to LF refining and RH refining.
  • LF refining the stirring time after all alloys are added is ⁇ 3 min.
  • RH refining the amount of calcium wire fed is 200-350 m
  • the argon blowing time after the calcium wire feeding is completed is ⁇ 5 min
  • the target Ca content is 10-40 ppm
  • the Ca/S ratio is ⁇ 2.
  • the heating temperature is 1100-1200°C and the holding time is 200-350 minutes;
  • the final rolling temperature is 750 ⁇ 900°C
  • the start cooling temperature is 700-800°C
  • the stop cooling temperature is 300-550°C
  • the cooling rate is ⁇ 30°C/s
  • the cooling rate is preferably 10-30°C/s.
  • Composition design and internal quality control are crucial to improving the acid resistance and toughness of steel.
  • the increase in S content in the process and finished product will increase the precipitation and aggregation of inclusions such as MnS in the steel.
  • MnS inclusions produced in the core are elongated after rolling and are prone to form gaps after cooling.
  • MnS is one of the most likely locations to form HIC, which greatly reduces low-temperature toughness and HIC resistance.
  • step 1) of the method of the present invention the S content of the molten iron before pretreatment is controlled to be ⁇ 0.050%, the S content of the molten iron entering the furnace is controlled to be ⁇ 0.0020%, and the molten iron ladle is cleaned after desulfurization.
  • step 2) of the method of the present invention during the smelting process, the bottom blowing flow rate of the converter bottom calming stage is ⁇ 900Nm 3 /h, thereby enhancing the stirring of the molten pool, making the composition and temperature more uniform, and improving the quality of the molten steel; the free oxygen is controlled to be ⁇ 900ppm when the blowing is stopped, ensuring that the molten steel has a lower oxygen content when the blowing is stopped, reducing defects such as pores and inclusions in the subsequent solidification process, and improving the quality of the slab.
  • the increase in C content will aggravate the segregation degree in the steel, causing MnS and other inclusions to further gather in the core, and the tendency of hydrogen-induced cracking will increase significantly. Therefore, the C and S of the ladle molten steel need to be strictly controlled.
  • the C content of the ladle molten steel i.e., tapping
  • the C content of the ladle molten steel is controlled to be C ⁇ 0.035% and S ⁇ 0.00
  • step 3 of the method of the present invention in LF refining, according to the pre-desulfurization and pre-deoxidation conditions after the converter, aluminum slag is used for slag deoxidation, and the purpose of desulfurization is achieved by stirring with argon gas at the bottom, and the stirring time is controlled to be ⁇ 3min.
  • the target wire feeding amount of silicon calcium wire in RH refining is 200-350m, and after the wire feeding of RH refining is completed, a small flow of argon is blown from the bottom of the ladle for more than 5min, and the target Ca is controlled to be 10-40ppm and the Ca/S ratio is ⁇ 2 , so as to modify the inclusions such as MnS and Al2O3 in the steel.
  • the slab heating temperature is 1100-1200°C; the slab insulation time is 200-350 minutes. Sufficient heating temperature and time are used to ensure that the alloy is fully dissolved and has good uniformity. It is beneficial to obtain uniform structure and plate shape.
  • the final rolling temperature is controlled to be 750-900°C.
  • the austenite grains are obviously coarsened, which will have an adverse effect on the final steel plate structure and performance, while when the final rolling temperature is too low, mixed crystal structure is easily generated in the steel, which has an adverse effect on the toughness, plasticity and processing properties of the steel.
  • the rolling adopts the TMCP process.
  • step 7) of the method of the present invention after rolling, a DQ+ACC cooling mode is adopted, and the addition amount of Mo and Nb elements is limited to a relatively low level, and their inhibitory effect on block ferrite is appropriately reduced, so that the steel undergoes phase transformation under the conditions of a start cooling temperature of 700-800°C, a stop cooling temperature of 300-550°C, and a cooling rate of ⁇ 30°C/s (preferably 10-30°C/s).
  • a small amount of ferrite is first precipitated along the original austenite grain boundary, and has sufficient phase transformation time to transform into block ferrite with a relatively regular shape, thereby obtaining a small amount of block ferrite ( ⁇ 10 volume %) without affecting the phase transformation of bainite.
  • block ferrite can increase the interface area and cause a large plastic deformation at the crack tip, relieve the triaxial tensile stress at the crack tip, and hinder crack propagation; at the same time, the presence of granular bainite inhibits the coordinated deformation of ferrite and increases deformation resistance. Most cracks need to frequently cross the interface between bainite and polygonal ferrite during the expansion process.
  • the crack development is effectively suppressed, thus showing good low-temperature DWTT toughness.
  • the content of blocky ferrite is 3-10% by volume.
  • the organization's ability to coordinate deformation is insufficient, the overall toughness and plasticity are poor, and stress concentration is easily generated, resulting in rapid crack propagation; when it is higher than 10% by volume, the organization's deformation resistance is low, the overall toughness is insufficient, and the ability to resist crack development is poor.
  • a cooling rate of 10 to 30 °C/s is conducive to rapid and uniform transformation of the organization, reducing core organization segregation and inclusion aggregation, and avoiding local hard point defects on the surface, obtaining uniformly refined organization grains in the thickness direction and throughout the plate, ensuring uniform and stable overall performance of the steel plate, and ensuring that the steel plate has excellent acid resistance while having good low-temperature toughness.
  • the method satisfies at least one of the following conditions:
  • step 2) the C content of the tapped steel is 0.018% to 0.035%, and/or the S content of the tapped steel is 0.0036% to 0.0080%;
  • step 4 the continuous casting calming time is 3 to 12 minutes;
  • step 7 the cooling rate is 15-30°C/s.
  • the present invention has the following advantages:
  • the present invention adopts ultra-low C, ultra-low S and low inclusion element solutions.
  • Nb+Ti ⁇ 0.07%, 0.7 ⁇ (Nb+27.5*Ti)/(5C+27*N) ⁇ 5.0 are controlled, and the influence of second phase particles such as TiN and NbC on the acid resistance and low-temperature toughness of steel is focused on.
  • the present invention adopts DQ+ACC cooling mode after rolling, controls the start cooling temperature to 700-800°C, the stop cooling temperature to 300-550°C, and the cooling rate ⁇ 30°C/s (preferably 10-30°C/s).
  • the inhibitory effect of elements such as Mo on blocky ferrite is reduced, a uniform and refined granular bainite + a small amount of blocky ferrite microstructure can be formed in this mode, wherein the blocky ferrite content is 3-10%, the grain size is ⁇ 10 levels, and there is no obvious segregation band or hard phase structure.
  • Blocky ferrite as a tough phase, can increase the interface area and cause greater plastic deformation at the crack tip, thereby alleviating the triaxial tensile stress at the crack tip and hindering crack propagation, thereby making the acid-resistant pipeline steel have good low-temperature toughness.
  • the pipeline steel of the present invention has better low-temperature toughness than the prior art, with a -49°C Charpy impact energy ⁇ 400 J, and a -35°C DWTT SA% ⁇ 90%.
  • the pipeline steel with the best low-temperature toughness is a -40°C Charpy impact energy ⁇ 300 J, and a -30°C DWTT SA% ⁇ 85%.
  • Figure 1 is a scatter plot of the HIC crack length rate CLR and the coefficient Ks. As shown in the figure, when Ks ⁇ 0.23, CLR is 0, that is, no cracking occurs.
  • FIG. 2 is a microstructure photograph of the pipeline steel according to Example 1 of the present invention.
  • the bottom blowing flow rate of the converter bottom calming stage is ⁇ 900Nm 3 /h
  • the free oxygen is controlled to be ⁇ 900ppm when the blowing is stopped
  • the C content of the tapped steel is controlled to be ⁇ 0.035%
  • the S content of the tapped steel is controlled to be ⁇ 0.0080%
  • the steel produced in step 2) is subjected to LF refining and RH refining.
  • LF refining the stirring time after all alloys are added is ⁇ 3 min.
  • RH refining the amount of calcium wire fed is 200-350 m
  • the argon blowing time after the calcium wire feeding is completed is ⁇ 5 min
  • the target Ca content is 10-40 ppm
  • the Ca/S ratio is ⁇ 2.
  • the heating temperature is 1100-1200°C and the holding time is 200-350 minutes;
  • the DQ+ACC cooling mode After rolling, the DQ+ACC cooling mode is adopted, the start cooling temperature is 700-800°C, the stop cooling temperature is 300-550°C, and the cooling rate is 10-30°C/s.
  • Example 1 The microstructure morphology of Example 1 was observed by an optical microscope (manufacturer: ZEISS, model: Axio Imager.M2m), and the volume percentage content and grain size of the bulk ferrite were measured by metallographic microscopy.
  • Hydrogen sulfide stress corrosion cracking (SSC) resistance test According to the NACE TM 0177-2016 standard "Test method for resistance of metallic materials to sulfide stress corrosion cracking in pressurized sulfur environments", the hydrogen sulfide stress corrosion cracking (SSC) resistance was measured by using a four-point bending method with a stress loading of 90% SYMS. After 720 hours of immersion in solution A, the measurement results are shown in Table 5. None of the samples obtained in the embodiments of the present invention were broken.
  • the Rt0.5 yield strength of the pipeline steel obtained by the present invention is 430-570 MPa
  • the tensile strength is 520-690 MPa
  • the yield ratio is ⁇ 0.92
  • the hardness HV10 ⁇ 220 and it has good low-temperature toughness, -49°C Charpy impact energy ⁇ 400 J, -35°C DWTT SA% ⁇ 90%
  • the content of massive ferrite is 3-10 volume %
  • the grain size is ⁇ 10 grade.
  • Figure 2 also shows a microstructure photograph of the pipeline steel obtained by Example 1 of the present invention.
  • the steel plate obtained by the chemical composition design and manufacturing method of the present invention is a microstructure of uniformly refined granular bainite and a small amount of massive ferrite; wherein the content of massive ferrite is 3-10%, the grain size is ⁇ 10 grade, and there is no obvious segregation band or hard phase structure.
  • the C content is 0.042 (outside the range of C of the present invention: 0.015-0.040), and Ks reaches 0.311 (Ks ⁇ 0.23 of the present invention), (Nb+27.5*Ti)/(5C+27*N) is 0.61 ( 0.7 ⁇ 5.0), Cu+Ni+Cr+Mo reaches 0.93% (Cu+Ni+Cr+Mo ⁇ 0.80% of the present invention), the comparative steel of Comparative Example 1 has large cracks in the HIC test, two of the three specimens are cracked, the crack length rate CLR exceeds 20%, cracks also appear in the SSC test, and the DWTT and impact properties are worse than those of the embodiments.
  • Ks reaches 0.289 (Ks ⁇ 0.23 of the present invention)
  • the length of the calcium-feeding wire is relatively small
  • the Ca/S ratio is only 1.0 (Ca/S ⁇ 2 of the present invention)
  • the B content reaches 0.0005% (B content ⁇ 0.0004% of the present invention)
  • O+N+H reaches 0.0072% (O+N+H ⁇ 0.0070% of the present invention).
  • the comparative steel of Comparative Example 2 has poor HIC resistance and toughness.
  • the crack length rate CLR of one sample reaches 18.56%
  • cracks also appear in the SSC test
  • the DWTT is only 70%/75%
  • the average impact energy is 279J
  • one of the single values is only 208J.
  • the pipeline steel of the present invention can achieve the excellent effect of having high strength and acid resistance as well as good low-temperature toughness through reasonable component matching and process design.

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Abstract

本发明涉及一种管线钢及其制造方法，该管线钢除含有Fe和不可避免的杂质之外，还含有按质量百分比计的如下化学元素：C 0.015～0.040％、Si 0.10～0.30％、Mn 1.18～1.45％、0＜P≤0.008％、0＜S≤0.0010％、Cu 0.05～0.35％、Ni 0.05～0.30％、Cr 0.15～0.30％、Mo 0.02～0.09％、Nb 0.025～0.055％、Ti 0.005～0.020％、Ca 0.0010～0.0040％、Alt 0.010～0.040％、0＜B≤0.0004％、0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％、0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％，并且满足：Nb+Ti≤0.07％；Ca/S≥2；Cu+Ni+Cr+Mo≤0.80％；0＜N+O+H≤0.0070％；0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0；氢致开裂敏感系数Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19。得到的管线钢能够达到X60/X65强度级别，并且具有良好的低温韧性以及优异的抗HIC、SSC性能。

Description

一种管线钢及其制造方法 技术领域
本发明属于金属材料及其制造领域，尤其涉及一种管线钢及其制造方法。
背景技术
随着常规油气田的不断开发，储量不断减少，含H2S酸性油气田的开发比例进一步增加，这就要求输送管线管具有抗酸性能。不同的抗酸管线工程的设计方会根据管道服役条件制定个性化的抗酸相关指标，包括NACE标准的抗氢致开裂(HIC)性能、抗硫化氢应力腐蚀开裂(SSC)性能，以及严格的低倍评级、探伤要求等。
中国专利CN202010552764.7公开了“一种耐酸性能优良的含镁X65管线钢及生产方法”，该管线钢组分及重量百分比含量为：C：0.03～0.05％、Si：0.17～0.20％、Mn：0.40～0.70％、Ni：0.12～0.15％、Cr：0.45～0.48％、Mo：0.10～0.15％、Mg：0.002～0.004％、Nb：0.06～0.09％、Ti：0.01～0.02％，P≤0.008％、S≤0.010％，其余为Fe及不可避免的杂质。该专利采用Mg处理方法进行夹杂物处理，化学成分中Mn含量较低，因此通过添加较多的Nb、Cr及一定量的Mo以保证强度，获得X65强度级别的抗酸管线钢。但是，使用Mg来处理夹杂物的方式对稳定添加进行控制的难度较高，并且对MnS等夹杂物的控制效果较差，对于抗酸及韧性均有不利影响，所得抗酸管线钢-20℃夏比冲击功≥200J，-15℃DWTT SA％≥90％。
中国专利CN201510870777.8公开了“一种厚规格抗酸性X60MS热轧卷板及其制造方法”，该钢板的成分按重量百分比计如下：C：0.050～0.070％、Si：0.15～0.25％、Mn：1.35～1.45％、Nb：0.025～0.045％、Ti：0.008～0.025％、Mo：0.06～0.10％、Cr：0.08～0.18％、Als：0.015～0.045％、P≤0.015％、S≤0.0015％、N≤0.006％，由于C含量增加会显著增加钢中的偏析程度，增加钢材的HIC敏感性，并在一定程度上降低钢材的低温韧性，所得钢材的-40℃夏比冲击功≥300J，-30℃DWTT SA％≥85％。
中国专利CN202011197059.6公开了“一种X60/X65级抗酸管线钢及其制备方法”，其化学成分及质量百分含量如下：C：0.02～0.06％，Si：0.10～0.20％，Mn≤0.90％，Alt：0.020～0.050％，Nb：0.060～0.080％，Ti≤0.020％，Cr≤0.30％，Ni≤0.20％，Cu≤0.40％，P≤0.012％，S≤0.002％，N≤0.0060％，B≤0.0005％，Mn+Cr≤1.2％，其余为铁及不可避 免的杂质元素。该专利合金成分的设计方案添加较多Nb(0.060～0.080％)，高Nb含量使得钢中含Nb颗粒在芯部偏聚的倾向性增加，钢板芯部出现HIC裂纹的概率随之增加。
从现有专利可发现，现有的抗酸管线钢中添加的C和Nb等易偏析元素含量过高，容易导致芯部出现偏析及夹杂，提高了钢材料的抗HIC敏感性，出现氢致开裂的可能性增大，且钢材的低温韧性难以进一步提升。
发明内容
本发明的目的在于提供一种高韧性X60/X65级抗酸管线钢及其制造方法，该管线钢能够达到X60/X65强度级别，并且具有良好的低温韧性以及优异的抗酸性能(例如抗HIC性能、抗SSC性能)，所述管线钢Rt0.5屈服强度为430～570MPa，抗拉强度为520～690MPa，屈强比≤0.92，延伸率A50≥30％，硬度HV10≤220，-49℃夏比冲击功≥400J，-35℃DWTT SA％≥90％。
为达到上述目的，一方面，本发明提供了一种管线钢，其除了含有Fe和不可避免的杂质之外，还含有按质量百分比计的如下化学元素：C：0.015～0.040％、Si：0.10～0.30％、Mn：1.18～1.45％、0＜P≤0.008％、0＜S≤0.0010％、Cu：0.05～0.35％、Ni：0.05～0.30％、Cr：0.15～0.30％、Mo：0.02～0.09％、Nb：0.025～0.055％、Ti：0.005～0.020％、Ca：0.0010～0.0040％、Alt：0.010～0.040％、0＜B≤0.0004％、0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％、0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％，并且满足：
Nb+Ti≤0.07％；
Ca/S≥2；
Cu+Ni+Cr+Mo≤0.80％；
0＜N+O+H≤0.0070％；
0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0；和
氢致开裂敏感系数Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19，式中元素符号代入其对应元素含量的质量百分号前的数值。
在本文中，“不可避免的杂质”指除P、S之外的其他不可避免的杂质。
优选地，所述管线钢含有按质量百分比计的如下化学元素：C：0.015～0.040、Si：0.10～0.30％、Mn：1.18～1.45％、0＜P≤0.008％、0＜S≤0.0010％、Cu：0.05～0.35％、Ni： 0.05～0.30％、Cr：0.15～0.30％、Mo：0.02～0.09％、Nb：0.025～0.055％、Ti：0.005～0.020％、Ca：0.0010～0.0040％、Alt：0.010～0.040％、0＜B≤0.0004％、0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％、0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％，余量为Fe和不可避免的杂质，并且满足：
Nb+Ti≤0.07％；
Ca/S≥2；
Cu+Ni+Cr+Mo≤0.80％；
0＜N+O+H≤0.0070％；
0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0；和
氢致开裂敏感系数Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19，式中元素符号代入其对应元素含量的质量百分号前的数值。
本发明所述的管线钢的化学成分设计原理如下：
C：C元素为钢中最经济的强化元素，其通过间隙固溶强化提升钢的强度。提高碳含量可大幅提升钢的淬透性，减少其他贵重合金元素添加量，降低生产成本，但C含量增加对钢的延性、韧性、焊接性和耐腐蚀性能不利，尤其容易加重芯部偏析，对抗酸性能有显著的不利影响。因此，本发明中采用超低C的设计，其质量百分比含量控制在0.015～0.040％，优选为0.017～0.032％。
Si：Si元素为固溶强化元素，同时也是钢中的脱氧元素，但Si含量过高会对表面质量及焊接性能产生不利影响，若Si含量超出0.30％，韧性可能下降。因此，本发明中将Si的质量百分比含量控制在0.10～0.30％，优选为0.18～0.27％。
Mn：Mn元素通过固溶强化提高钢的强度，是钢中补偿因C含量降低而引起强度损失的最主要且最经济的强化元素。Mn有助于获得细小的相变产物，也有助于在炼钢过程中对氧和硫的控制。如果Mn含量过低，将难以达到目标强度水平，但Mn含量过高也会加重中心偏析。因此，本发明中将Mn的质量百分比含量控制在1.18～1.45％，优选为1.25～1.43％,。
Cr：Cr元素具有一定的固溶强化作用，并可有效提升钢的淬透性。Cr含量在0.15％以上时，能有效改善钢的耐腐蚀性能，在钢的表面形成较为致密的保护层，起到保护基体的作用。但当钢中的Cr含量过高时，将不利于焊缝质量，易形成灰斑缺陷。因此，本发明中将Cr的质量百分比含量控制在0.15～0.30％，优选为0.20～0.25％。
Nb：Nb元素为低碳微合金钢的重要元素，热轧过程中固溶的Nb经应变诱导析出 从而形成Nb的碳氮化物，钉扎晶界抑制形变奥氏体的长大，经控制轧制和控制冷却使形变奥氏体相变为具有高位错密度的细小产物。固溶Nb在钢带卷取后会以第二相粒子NbC在基体内弥散析出，起到析出强化的作用。过多的Nb使得板坯容易产生裂纹，从而影响表面质量，此外还会恶化焊接性能。因此，本发明中将Nb的质量百分比含量控制在0.025～0.055％，优选为0.035～0.045％。
Ti：Ti元素为良好的脱氧去气剂和固定氮碳的有效元素。Ti的未溶碳氮化物在钢加热时可以阻止奥氏体晶粒的长大，在高温奥氏体区粗轧时析出的TiN和TiC可有效抑制奥氏体晶粒长大从而细化晶粒，同时还能提高Nb的固溶度，降低含Nb钢的微裂纹敏感性，Ti一般与Nb复合加入，另外在焊接过程中析出也能抑制高温晶粒长大从而改善焊接性能。因此，本发明中将Ti的质量百分比含量控制在0.005～0.020％，优选为0.007～0.018％。
Nb+Ti：添加过多的Nb、Ti容易导致钢板中出现较粗的含Nb/Ti颗粒的聚集，容易在芯部形成沿轧制方向分布的长条状偏聚，该长条状偏聚为硬质相并且其硬度高于基体，增加组织的氢致开裂敏感性，使氢钉扎在其周围不再扩散，因此容易成为氢聚集处，也易成为氢致开裂的起始位置，显著增加芯部HIC开裂倾向，恶化钢材的抗酸性能，因此，本发明中控制Nb、Ti元素的的质量百分比含量满足Nb+Ti≤0.07％，优选为0.033～0.069％，更优选为0.033～0.065％。
另外，TiN、NbC是钢中重要的第二相，对于强韧性等具有重要影响。当(Nb+27.5*Ti)/(5C+27*N)＞5时，第二相颗粒较为粗大、分布较为不均，易在芯部偏析处聚集，造成应力集中，引发脆性断裂，导致钢材的低温韧性急剧减小，显著降低钢材及焊缝热区韧性，还会成为导致形成氢致开裂的硬质相。当(Nb+27.5*Ti)/(5C+27*N)＜0.7时，析出相的数量过少，且尺寸过于细小，容易再固溶进钢材中，晶界钉扎效果微弱，同样对韧性不利。因此，本发明中控制Nb、Ti元素的的质量百分比含量满足0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0，优选满足0.81≤(Nb+27.5*Ti)/(5C+27*N)≤3.69，更优选1.13≤(Nb+27.5*Ti)/(5C+27*N)≤2.70。
Mo：Mo为强淬透性元素，其能够显著推迟铁素体相变，抑制铁素体和珠光体的形成，有效促进贝氏体转变并起到强化基体的作用，得到更加细小的组织。一定量的Mo对钢材的组织性能改善有明显益处，但当Mo含量过高时，钢的塑性降低，并且Mo的成本高昂。因此，本发明中将Mo的质量百分比含量控制在0.02～0.09％的低水平。
Cu、Ni：Cu元素和Ni元素可通过固溶强化作用提高钢的强度，并改善钢的耐大 气腐蚀性能，但Cu元素和Ni元素含量过高时会导致钢板热裂纹，而Ni则可细化晶粒，并改善Cu在钢中易引起的热脆性，对于低温韧性提升有显著帮助。因此，本发明中将Cu的质量百分比含量控制在0.05～0.35％，优选为0.15～0.33％；将Ni的质量百分比含量控制在0.05～0.30％，优选为0.10～0.22％。
Cu+Ni+Cr+Mo：Cu、Ni、Cr、Mo元素的添加会增加生产成本，并且如果Cu、Ni、Cr、Mo的含量过高，将会增加钢材的热脆性倾向，同时会促进MA颗粒形成，增加钢中的硬质相比例，从而成为氢陷阱，产生氢脆性、脆性相等系列问题，从而降低抗HIC性能及低温韧性，影响钢材的性能。为确保钢材的强度、韧性等综合性能，并降低生产成本，本发明中控制Cu、Ni、Cr、Mo元素的质量百分比含量满足Cu+Ni+Cr+Mo≤0.80％，优选满足0.34≤Cu+Ni+Cr+Mo≤0.79％。
S、P：S、P元素是钢中主要的杂质元素，P元素易导致钢的冷脆，S元素易引起热脆，从而导致钢的性能不稳定，尤其是随着S含量提高，MnS夹杂物增多，会显著增加钢的氢致开裂敏感性，降低钢的抗酸性能，因此，应尽量降低钢中的磷、硫的含量。因此，本发明中控制S、P元素的质量百分含量为S≤0.0010％，P≤0.0080％。
Ca：通过Ca处理可以控制硫化物的形态，改善钢板的各向异性，提高低温韧性；另外，当S含量过高时，控制适当的Ca/S比以确保钢材的抗HIC性能，但Ca含量过高容易导致钢中CaO等夹杂物增多增大。本发明中将Ca的质量百分比含量控制在0.0010～0.0040％，优选为0.0021～0.0035％，同时控制Ca/S≥2，优选地，Ca/S比为3～13，更优选为4.1～8.7。
Al：Al元素用于钢的脱氧，适量Al元素还有利于细化晶粒，改善强韧性能，但若总铝(Alt)含量大于0.04％，则可能形成粗的沉淀物从而降低钢的抗酸性能及低温韧性。因此，本发明将Alt的质量百分比含量控制在0.010～0.040％。
N、O、H：对于低温韧性，O、N和H元素都可能导致脆性断裂的发生，尤其是在低温下。O和N的存在会形成O化物和N化物，这些物质可能在晶界处聚集，导致晶间脆性的发生。同时，H也可能会渗入晶粒中，使得晶格变形、形成孔洞，从而导致脆性断裂。因此，钢的低温韧性与O、N、H的含量有关，过高或不当的含量都可能导致脆性断裂的发生。对于抗酸性能，O、N和H的影响也比较显著。在酸性环境下，O会与钢中的铁元素发生氧化反应，使钢材失去一定的抗腐蚀能力，N则可能与钢中的铬元素形成N化物，导致钢材的耐腐蚀性下降，而H则会促进钢的应力腐蚀开裂，对钢的抗酸性也会造成一定的影响。因此，本发明中将N、O、H元素的质量百分比含量控制为0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％，并且0＜ N+O+H≤0.0070％，优选为0.0024％≤N+O+H≤0.0065％。
B：B元素在钢中的主要作用是增加钢的淬透性和强度，从而节约其他稀有贵重的金属，但B元素的加入会对钢材的低温韧性产生显著的不利影响。因此，本发明中将B元素的质量百分比含量控制为0＜B≤0.0004％。
Sn、Sb、Bi、Pb、As：Sn、Sb、Bi、Pb、As元素在钢中的含量超过一定限度时，会降低钢的强度，并且容易造成严重偏析，在晶界形成偏聚，导致脆性裂纹从晶界开始，沿晶界延伸直至完全断裂，显著降低钢材的低温韧性；此外在钢中Sn含量过高会显著降低钢的抗HIC氢致开裂性能，这是因为Sn会促进氢的扩散和聚集，从而加速氢致开裂的发生，Pb也具有类似的作用。因此，本发明中将Sn、Sb、Bi、Pb、As元素的质量百分比含量控制为0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％。
氢致开裂敏感系数Ks：抗HIC性能受钢中各种元素的影响，如其中偏析程度较显著的元素C、S，参与第二相颗粒析出的Nb、Ti，以及N、B等夹杂元素，根据各元素的易偏析程度，以及对各元素与抗HIC性能之间的规律的研究，本发明中提出氢致开裂敏感系数Ks，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19，其中各元素符号表示各元素含量的质量百分号前的数值。HIC裂纹长度率CLR与系数Ks的相关性如图1所示，随着Ks增加，钢材的氢致开裂敏感性越高，出现HIC裂纹的概率越高，裂纹的长度也越长。当Ks≤0.23时，经受HIC试验的钢材均未开裂，当Ks＞0.23时，CLR开始显著升高。API 5L SPEC标准对于CLR的要求为≤15％，另外一些工程项目要求CLR≤10％甚至更加严格。根据CLR与Ks的关系，为保证钢材的抗酸性能，本发明中控制Ks≤0.23。
总体上，本发明采用上述设计的成分体系，综合考虑了钢中C、S等易偏析元素，Cu、Ni、Cr、Mo和Nb、Ti微合金化元素，以及N、O、H、B和Sn、Sb、Bi、Pb、As等元素对钢材的抗酸性能和低温韧性的影响，并根据各元素的易偏析程度以及与抗HIC性能的关系，提出对氢致开裂敏感系数Ks的控制要求，以提高钢材的抗酸性能。在此基础之上，重点考虑TiN、NbC等第二相颗粒对钢材抗酸性能和低温韧性的影响，控制Nb+Ti≤0.07％，0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0，同时，通过控制Cu+Ni+Cr+Mo≤0.80％，降低对块状铁素体的抑制作用，从而使钢组织中含有少量块状铁素体(块状铁素体含量3～10体积％)，由此抑制裂纹的产生，从而使管线钢在具有较高强度和抗酸性能的基础上，还具有较好的低温韧性。
优选地，所述管线钢的微观组织为粒状贝氏体+块状铁素体。
优选地，所述块状铁素体含量为3～10体积％，更优选为4.6～8.5体积％。
优选地，所述块状铁素体的晶粒度≥10级，更优选为10.5～11级。
优选地，所述粒状贝氏体的含量为90～97体积％，更优选为91.5～95.4体积％。
优选地，所述粒状贝氏体为均匀细化的粒状贝氏体。
优选地，所述管线钢的微观组织中没有明显的偏析带或硬相组织。
优选地，所述管线钢的Rt0.5屈服强度为430MPa以上，优选为450～570MPa，更优选为510～570MPa；抗拉强度为520MPa以上，优选为540～690MPa，更优选为615～690MPa；屈强比≤0.92，优选为≤0.87，更优选为0.81～0.87；延伸率A50≥30％，优选为30％～47％，更优选为40％～47％；硬度HV10≤220，优选为195～220；-49℃夏比冲击功≥400J，优选≥430J，更优选为430J～500J；-35℃DWTT SA％≥90％，优选为90～98％。
优选地，所述管线钢满足如下性能中的至少一项：在根据NACE TM0284-2016标准《管道和压力容器用钢抗氢致开裂性能评价的试验方法》进行的抗HIC性能测试中无裂纹；在根据NACE TM 0177-2016标准《金属材料在含压S环境中抗硫化物应力腐蚀开裂性能试验方法》进行的抗SSC性能测试中无开裂。
另一方面，本发明还提供了制造上述管线钢的方法，其包括如下步骤：
1)铁水预处理
控制预处理前铁水的S含量≤0.050％，入炉铁水的S含量≤0.0020％；
2)冶炼
对步骤1)的铁水进行冶炼，控制出钢的C含量≤0.035％，出钢的S含量≤0.0080％；
3)精炼
对步骤2)的出钢进行LF精炼和RH精炼，在LF精炼期间，合金全部添加完后的搅拌时间≥3min，在RH精炼期间，喂钙线的量为200～350m，喂钙线结束后的吹氩时间≥5min，目标Ca含量为10～40ppm，Ca/S比≥2；
4)连铸
连铸镇静时间≥3min，过热度控制在10～40℃；
5)再加热
加热温度为1100～1200℃，保温时间为200～350min；
6)控制轧制
终轧温度为750～900℃；
7)轧后冷却
开冷温度为700～800℃，停冷温度为300～550℃，冷速≤30℃/s，冷速优选为10～30℃/s。
优选地，在步骤2)中，转炉底镇静阶段底吹流量≥900Nm3/h，控制停吹游离氧≤900ppm。
成分设计及对内部质量的控制对于改善钢材的抗酸性能及韧性至关重要。过程以及成品的S含量增多会增加钢中MnS等夹杂物的析出聚集，芯部产生的MnS夹杂物经轧制被拉长，冷却后易形成间隙，MnS是最易形成HIC的位置之一，大幅降低低温韧性和抗HIC性能。为保证成品及过程中含有较低的S，在本发明方法的步骤1)中控制预处理前的铁水的S含量≤0.050％，控制入炉铁水的S含量≤0.0020％，脱硫后铁水包扒渣干净。
在本发明方法的步骤2)中，在冶炼过程中，转炉底镇静阶段底吹流量≥900Nm3/h，从而增强熔池搅拌，使成分及温度更加均匀，改善钢水质量；控制停吹游离氧≤900ppm，确保停吹时钢水具有较低氧含量，减少后续凝固过程中气孔、夹杂等缺陷，提高板坯质量。C含量升高会加剧钢中偏析程度，使得MnS等夹杂在芯部进一步聚集，发生氢致开裂的倾向会显著增加，因此，对于钢包钢水的C和S需进行严格控制，本发明中钢包钢水(即出钢)的C含量控制为C≤0.035％，S≤0.0080％。
在本发明方法的步骤3)中，在LF精炼中，根据转炉炉后预脱硫和预脱氧情况，渣脱氧采用铝渣，通过底搅氩气搅拌达到脱硫的目的，控制搅拌时间≥3min。RH精炼中硅钙丝目标喂丝量为200～350m，RH精炼喂丝结束后进行小流量钢包底吹氩5min以上，控制目标Ca为10～40ppm、Ca/S比≥2，以对钢中的MnS和Al2O3等夹杂物进行改性。通过增加钢中有效钙含量，在钢水凝固过程中提前形成高熔点CaS，抑制此过程中生成MnS的总量，并将其全部或部分改性成CaS，另一方面使大颗粒Al2O3夹杂物变成低熔点复合夹杂物，促进其上浮，净化钢水。LF及RH全部使用低碳合金，即C含量控制在0.015～0.040％。
在本发明方法的步骤4)中，在连铸过程中，控制连铸镇静时间≥3分钟，以减少板坯缩孔、疏松等缺陷，提高板坯质量。过热度较高会加重偏析程度，过低则夹杂物无法充分上浮，均对抗酸及低温韧性产生不利影响。因此，本发明中将过热度控制在10～40℃，以控制中心偏析，低倍评级曼内斯曼标准≤2级。优选地，采用全程保护、恒速浇注的方式确保板坯具有良好内部质量及偏析水平。
在本发明方法的步骤5)中，板坯加热温度为1100～1200℃；板坯保温时间为200～350min。通过足够的加热温度和时间保证合金充分固溶，并且有较好的均匀性， 有利于获得均匀的组织及板形。
在本发明方法的步骤6)中，终轧温度控制为750～900℃。当终轧温度过高时，奥氏体晶粒粗化明显，这将对最终钢板组织性能产生不利影响，而当终轧温度偏低时，钢中易产生混晶组织，对钢材的韧性、塑性及加工性能产生不利影响。优选地，轧制采用TMCP工艺。
在本发明方法的步骤7)中，轧后采用DQ+ACC冷却模式，通过将Mo、Nb元素的添加量限制在较低水平，适当降低其对块状铁素体的抑制作用，使得钢在开冷温度为700～800℃，停冷温度为300～550℃，冷速≤30℃/s(优选10～30℃/s)的条件下进行相变，此时少量铁素体首先沿原奥氏体晶界析出，并有足够的相变时间转变为形状相对规则的块状铁素体，从而得到少量块状铁素体(≤10体积％)同时不影响贝氏体的相变。块状铁素体作为韧性相，可以增加界面面积，并使裂纹尖端发生较大塑性变形，缓解裂纹尖端的三向拉应力，阻碍裂纹扩展；同时粒状贝氏体的存在抑制了铁素体的协调变形，增加了变形抗力。大部分裂纹在扩展过程中需要频繁穿过贝氏体和多边形铁素体的界面，当组织中具有少量比例的块状形铁素体时，裂纹发展被有效抑制，因而表现出较好的低温DWTT韧性。优选地，块状铁素体的含量为3～10体积％，当低于3体积％时，组织协调变形能力不足，整体韧塑性较差，容易产生应力集中，导致裂纹快速传播；当高于10体积％时，组织变形抗力较低，整体强韧性不足，抵抗裂纹发展的能力较差。
此外，10～30℃/s的冷速有利于组织快速均匀转变，减少芯部组织偏析及夹杂物聚集，并避免表面局部出现硬点缺陷，获得厚度方向及全板均匀细化的组织晶粒，确保钢板整体性能均匀稳定，确保钢板在具有良好的低温韧性的前提下，还能具有优异的抗酸性能。
优选地，所述方法满足如下条件中的至少一项：
在步骤2)中，出钢的C含量为0.018％～0.035％，和/或出钢的S含量为0.0036～0.0080％；
在步骤3)中，合金全部添加完后的搅拌时间为3～12min，喂钙线结束后的吹氩时间为5～14min，和/或Ca/S比为3～13；
在步骤4)中，连铸镇静时间为3～12min；
在步骤7)中，冷速为15～30℃/s。
与现有技术相比，本发明具有如下优点：
在成分设计方面，本发明采用超低C、超低S及低夹杂元素方案，根据各元素的 易偏析程度及与抗HIC性能关系，提出对氢致开裂敏感系数Ks的控制要求(即Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19)，以提高钢材的抗酸性能。在此基础之上，控制Nb+Ti≤0.07％，0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0，重点考虑TiN、NbC等第二相颗粒对钢材抗酸性能和低温韧性的影响，同时，通过控制Cu+Ni+Cr+Mo≤0.80％，降低对块状铁素体的抑制作用，从而使钢组织中含有少量块状铁素体(块状铁素体含量3～10％)，抑制裂纹的产生，从而使管线钢在具有较高强度和抗酸性能的基础上，还具有较好的低温韧性。而常规的抗酸管线钢中C、Mo、Nb元素含量较高，也没有考虑对夹杂元素(例如N、B等)的控制。
在制造方法方面，本发明在轧后采用DQ+ACC冷却模式，控制开冷温度为700～800℃，停冷温度为300～550℃，冷速≤30℃/s(优选10～30℃/s)，同时由于降低了Mo等元素对块状铁素体的抑制作用，在该模式下能够形成均匀细化的粒状贝氏体+少量块状铁素体的微观组织，其中块状铁素体含量3～10％，晶粒度≥10级，没有明显的偏析带或硬相组织，块状铁素体作为韧性相，可以增加界面面积，并使裂纹尖端发生较大塑性变形，缓解裂纹尖端的三向拉应力，阻碍裂纹扩展，从而是抗酸管线钢具有良好的低温韧性。
本发明的管线钢的Rt0.5屈服强度为430～570MPa，抗拉强度为520～690MPa，屈强比≤0.92，延伸率A50≥30％，硬度HV10≤220，-49℃夏比冲击功≥400J，-35℃DWTT SA％≥90％，具有优异的抗酸性能(在国际NACE标准A溶液下进行的抗HIC及SSCC试验(90％ SYMS)中均未开裂)，完全满足API 5L的评定标准。本发明的管线钢相对于现有技术具备更好的低温韧性-49℃夏比冲击功≥400J，-35℃DWTT SA％≥90％。而常规抗酸管线钢中，低温韧性最好的管线钢为-40℃夏比冲击功≥300J，-30℃DWTT SA％≥85％。
附图说明
图1为HIC裂纹长度率CLR与系数Ks的散点关系图。如图中所示，当Ks≤0.23时，CLR为0，即不发生开裂。
图2为本发明实施例1的管线钢的微观组织照片。
具体实施方式
下面结合实施例和附图对本发明的管线钢及其制造方法做进一步说明，该说明并 不对本发明的技术方案构成限定。
本发明实施例1-10的管线钢和对比例1-2的对比钢材均通过如下步骤制得：
1)铁水预处理
控制预处理前铁水的S含量≤0.050％，入炉铁水的S含量≤0.0020％；
2)冶炼
对铁水进行冶炼，转炉底镇静阶段底吹流量≥900Nm3/h，控制停吹游离氧≤900ppm，控制出钢的C含量≤0.035％，出钢的S含量≤0.0080％；
3)精炼
对步骤2)的出钢进行LF精炼和RH精炼，在LF精炼期间，合金全部添加完后的搅拌时间≥3min，在RH精炼期间，喂钙线的量为200～350m，喂钙线结束后的吹氩时间≥5min，目标Ca含量为10～40ppm，Ca/S比≥2；
4)连铸
连铸镇静时间≥3min，过热度控制在10～40℃；
5)再加热
加热温度为1100～1200℃，保温时间为200～350min；
6)控制轧制
终轧温度：750～900℃；
7)轧后冷却
轧后采用DQ+ACC冷却模式，开冷温度为700～800℃，停冷温度为300～550℃，冷速10～30℃/s。
通过上述步骤制得的本发明实施例1-10的管线钢和对比例1-2的对比钢材的化学元素成分如表1和表2所示，余量包括Fe及不可避免的杂质；本发明实施例1-10的管线钢和对比例1-2的对比钢材的制造方法的工艺参数如表3所示。实施例1-10的管线钢的化学元素成分和相关工艺设计均满足本发明设计要求；对比例1-2的对比钢材采用与本发明实施例1-10基本相同的方法制备，所不同的是其化学成分和/或工艺参数中的至少一个不符合本发明要求。
对制得的实施例1-10的管线钢和对比例1-2的对比钢材分别取样进行各项性能测试，所得的力学性能及微观组织测试结果如表4所示，抗HIC性能及SSC性能测试结果如表5所示。
相关性能测试的方法如下所述：
(1)拉伸、硬度、夏比冲击试验：根据ASTM A370进行拉伸、硬度、夏比冲击 试验，测试获得各实施例的管线钢和各对比例的对比钢材在室温下的Rt0.5屈服强度、抗拉强度、屈强比、延伸率A50，HV10维氏硬度，以及-49℃夏比冲击功数值。其中，屈强比＝屈服强度/抗拉强度。
(2)DWTT落锤撕裂试验：根据API RP 5L3测试获得的各实施例的管线钢和各对比例的对比钢材的-35℃DWTT SA％。
(3)微观组织观测：通过光学显微镜(生产商ZEISS，型号：Axio Imager.M2m)观察实施例1的微观组织形貌，并通过金相显微镜法测得块状铁素体的体积百分比含量和晶粒度。
(4)抗氢致开裂(HIC)性能测试：采用NACE TM0284-2016标准《管道和压力容器用钢抗氢致开裂性能评价的试验方法》，测量湿硫化氢(H2S)环境中抗氢致开裂(HIC)性能，将所得钢分别在标准试样及A溶液中持续浸泡96h，测量结果如表5所示，由本发明实施例1-10获得的试样均未出现裂纹。
(5)抗硫化氢应力腐蚀开裂(SSC)性能测试：依据NACE TM 0177-2016标准《金属材料在含压S环境中抗硫化物应力腐蚀开裂性能试验方法》，测量抗硫化氢应力腐蚀开裂(SSC)性能，采用四点弯曲法，应力加载为90％ SYMS，经过720小时A溶液浸泡，测量结果如表5所示，本发明获得的实施例试样均未断裂。



表5

由表4可知，本发明获得管线钢的Rt0.5屈服强度为430～570MPa，抗拉强度为520～690MPa，屈强比≤0.92，延伸率A50≥30％，硬度HV10≤220，并具有良好的低温韧性，-49℃夏比冲击功≥400J，-35℃DWTT SA％≥90％；块状铁素体含量为3～10体积％，晶粒度≥10级，此外，图2还显示了由本发明实施例1获得的管线钢的微观组织照片，由图2可以看出，采用本发明的化学成分设计以及制造方法获得的钢板为均匀细化的粒状贝氏体及少量块状铁素体的微观组织；其中，块状铁素体含量为3～10％，晶粒度≥10级，没有明显的偏析带或硬相组织。
由表5的测试结果可知，实施例1-10的管线钢抗氢致开裂(HIC)性能测试和抗硫化氢应力腐蚀开裂(SSC)性能测试中均未出现开裂。
在对比例1的对比钢材中，C含量为0.042(超出了本发明的范围C：0.015～0.040)，并且Ks达到0.311(本发明为Ks≤0.23)，(Nb+27.5*Ti)/(5C+27*N)为0.61(本发明为 0.7～5.0)，Cu+Ni+Cr+Mo达到0.93％(本发明的Cu+Ni+Cr+Mo≤0.80％)，对比例1的对比钢在HIC试验中出现较大裂纹，三个试样中有两个均开裂，裂纹长度率CLR超过20％，SSC试验同样出现裂纹，DWTT及冲击性能与实施例相比均更差。
在对比例2的对比钢材中，Ks达到0.289(本发明的Ks≤0.23)，喂钙丝长度偏小，Ca/S比仅为1.0(本发明的Ca/S≥2)，此外，B含量达到0.0005％(本发明的B含量≤0.0004％)，O+N+H达到0.0072％(本发明的O+N+H≤0.0070％)，对比例2的对比钢材抗HIC性能及韧性均较差，一个试样的裂纹长度率CLR达到18.56％，SSC试验也出现裂纹，DWTT仅为70％/75％，冲击功均值为279J，并且其中一个单值仅为208J。
综上所述，本发明的管线钢通过合理的成分匹配及工艺设计，能够实现在具有较高强度和抗酸性能的基础同时还具有良好的低温韧性的优异效果。
需要说明的是，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. 一种管线钢，其除含有Fe和不可避免的杂质之外，还含有按质量百分比计的如下化学元素：C：0.015～0.040％、Si：0.10～0.30％、Mn：1.18～1.45％、0＜P≤0.008％、0＜S≤0.0010％、Cu：0.05～0.35％、Ni：0.05～0.30％、Cr：0.15～0.30％、Mo：0.02～0.09％、Nb：0.025～0.055％、Ti：0.005～0.020％、Ca：0.0010～0.0040％、Alt：0.010～0.040％、0＜B≤0.0004％、0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％、0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％，并且满足：
    Nb+Ti≤0.07％；
    Ca/S≥2；
    Cu+Ni+Cr+Mo≤0.80％；
    0＜N+O+H≤0.0070％；
    0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0；
    氢致开裂敏感系数Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19，式中元素符号代入其对应元素含量的质量百分号前的数值。
  2. 如权利要求1所述的管线钢，其中，所述管线钢含有按质量百分比计的如下化学元素：C：0.015～0.040％、Si：0.10～0.30％、Mn：1.18～1.45％、0＜P≤0.008％、0＜S≤0.0010％、Cu：0.05～0.35％、Ni：0.05～0.30％、Cr：0.15～0.30％、Mo：0.02～0.09％、Nb：0.025～0.055％、Ti：0.005～0.020％、Ca：0.0010～0.0040％、Alt：0.010～0.040％、0＜B≤0.0004％、0＜O≤0.0026％、0＜N≤0.0050％、0＜H≤0.0002％、0＜Sn≤0.0050％、0＜Sb≤0.0025％、0＜Bi≤0.0030％、0＜Pb≤0.0050％、0＜As≤0.0030％，余量为Fe和不可避免的杂质，并且满足：
    Nb+Ti≤0.07％；
    Ca/S≥2；
    Cu+Ni+Cr+Mo≤0.80％；
    0＜N+O+H≤0.0070％；
    0.7≤(Nb+27.5*Ti)/(5C+27*N)≤5.0；
    氢致开裂敏感系数Ks≤0.23，Ks＝8.0C+24S-Alt-2.3Nb+2.8Ti+117B+0.34Ni-0.30Mo+24.5N-0.19，式中元素符号代入其对应元素含量的质量百分号前的数值。
  3. 如权利要求1或2所述的管线钢，其中，所述管线钢的微观组织为粒状贝氏体 +块状铁素体。
  4. 如权利要求3所述的管线钢，其中，所述块状铁素体的含量为3～10体积％，优选为4.6～8.5体积％。
  5. 如权利要求3或4所述的管线钢，其中，所述管线钢的晶粒度≥10级，优选为10.5～11级。
  6. 如权利要求1-5中任一项所述的管线钢，其中，所述管线钢的Rt0.5屈服强度为430MPa以上，优选为450～570MPa；抗拉强度为520MPa以上，优选为540～690MPa；屈强比≤0.92，延伸率A50≥30％，优选为≥40％；硬度HV10≤220，-49℃夏比冲击功≥400J，-35℃DWTT SA％≥90％。
  7. 如权利要求1-6中任一项所述的管线钢，其中，所述管线钢满足如下性能中的至少一项：在根据NACE TM0284-2016标准《管道和压力容器用钢抗氢致开裂性能评价的试验方法》进行的抗HIC性能测试中无裂纹；在根据NACE TM 0177-2016标准《金属材料在含压S环境中抗硫化物应力腐蚀开裂性能试验方法》进行的抗SSC性能测试中无开裂。
  8. 一种制造权利要求1～7中任一项所述的管线钢的方法，其包括如下步骤：
    1)铁水预处理
    控制预处理前铁水的S含量≤0.050％，入炉铁水的S含量≤0.0020％；
    2)冶炼
    对步骤1)的铁水进行冶炼，控制出钢的C含量≤0.035％，出钢的S含量≤0.0080％；
    3)精炼
    对步骤2)的出钢进行LF精炼和RH精炼，在LF精炼期间，合金全部添加完后的搅拌时间≥3min，在RH精炼期间，喂钙线的量为200～350m，喂钙线结束后的吹氩时间≥5min，目标Ca含量为10～40ppm，Ca/S比≥2；
    4)连铸
    连铸镇静时间≥3min，过热度控制为10～40℃；
    5)再加热
    加热温度为1100～1200℃，保温时间为200～350min；
    6)轧制
    终轧温度为750～900℃；
    7)轧后冷却
    开冷温度为700～800℃，停冷温度为300～550℃，冷速≤30℃/s。
  9. 如权利要求8所述的制造方法，其中，在步骤2)中，转炉底镇静阶段底吹流量≥900Nm3/h，控制停吹游离氧≤900ppm。
  10. 如权利要求8所述的制造方法，其中，所述方法满足如下条件中的至少一项：
    在步骤2)中，出钢的C含量为0.018％～0.035％，和/或出钢的S含量为0.0036～0.0080％；
    在步骤3)中，合金全部添加完后的搅拌时间为3～12min，喂钙线结束后的吹氩时间为5～14min，和/或Ca/S比为3～13；
    在步骤4)中，连铸镇静时间为3～12min；
    在步骤7)中，冷速为15～30℃/s。
PCT/CN2024/102421 2023-06-29 2024-06-28 一种管线钢及其制造方法 Ceased WO2025002365A1 (zh)

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