WO2014155440A1 - High strength thick steel plate for high heat input welding with excellent brittle crack arrestability and manufacturing method therefor - Google Patents

High strength thick steel plate for high heat input welding with excellent brittle crack arrestability and manufacturing method therefor Download PDF

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WO2014155440A1
WO2014155440A1 PCT/JP2013/006309 JP2013006309W WO2014155440A1 WO 2014155440 A1 WO2014155440 A1 WO 2014155440A1 JP 2013006309 W JP2013006309 W JP 2013006309W WO 2014155440 A1 WO2014155440 A1 WO 2014155440A1
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plate thickness
rolling
less
steel
brittle crack
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PCT/JP2013/006309
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French (fr)
Japanese (ja)
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長谷　和邦
佳子 竹内
三田尾　眞司
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Ｊｆｅスチール株式会社
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Priority to JP2013549634A priority Critical patent/JP5598618B1/en
Priority to CN201380075070.7A priority patent/CN105102650B/en
Priority to BR112015020815-0A priority patent/BR112015020815B1/en
Priority to KR1020157028785A priority patent/KR101732997B1/en
Publication of WO2014155440A1 publication Critical patent/WO2014155440A1/en
Priority to PH12015501719A priority patent/PH12015501719A1/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • 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
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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    • 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
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium 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
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—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
    • C22C38/48—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
    • 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
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • 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
    • 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
    • C21D2201/00—Treatment for obtaining particular effects
    • C21D2201/05—Grain orientation
    • 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

Definitions

  • the present invention relates to a high-strength thick steel plate for high-heat input welding having excellent brittle cracking arrestability and a method for producing the same, and more particularly to a ship. It relates to a plate having a thickness of 50 mm or more suitable for use.
  • Ni steel As a means of improving the brittle crack propagation stopping characteristics of steel materials, a method of increasing the Ni content has been conventionally known. In a liquefied natural gas (LNG) storage tank, 9% Ni steel is commercially available. Used on a scale.
  • LNG liquefied natural gas
  • TMCP Thermo-Mechanical Control Process
  • Patent Document 1 a steel material in which the structure of the surface layer portion is ultrafine (ultra-fine-grained-steel) is proposed in Patent Document 1.
  • Patent Document 1 focuses on the fact that shear lips (plastic deformation regions shear-lips) that occur in the steel surface layer when brittle cracks propagate are effective in improving the brittle crack propagation stopping characteristics.
  • shear lips plastic deformation regions shear-lips
  • Patent Document 1 discloses that the surface layer portion is cooled below the Ar3 transformation point by controlled cooling after hot rolling, and then the controlled cooling is stopped to bring the surface layer portion above the transformation point.
  • the process of recuperate is repeated one or more times, and during this time, the steel material is subjected to reduction, and it is repeatedly transformed or processed and recrystallized, so that a superfine ferrite structure or bainite structure is formed on the surface layer portion. (bainite structure) is generated.
  • both surface portions of the steel material have a circle-equivalent average grain. size): 5 ⁇ m or less
  • aspect ratio of aspect ratio: a layer having 50% or more of a ferrite structure having two or more ferrite grains, and suppressing variation in ferrite grain size is important.
  • the maximum rolling reduction per pass during finish rolling is set to 12% or less to suppress the local recrystallization phenomenon.
  • Patent Document 3 attention is paid not only to the refinement of ferrite crystal grains but also to subgrains formed in ferrite crystal grains, and a technique on the extension of TMCP that improves brittle crack propagation stop characteristics. Is described.
  • a) rolling conditions for securing fine ferrite crystal grains without requiring complicated temperature control such as cooling and recuperation of the steel sheet surface layer (b) Rolling conditions for generating a fine ferrite structure in a portion of 5% or more of the steel sheet thickness, (c) Dislocation introduced by machining (rolling) and development of texture in the fine ferrite by thermal energy
  • the brittle crack propagation stop property is improved by rolling conditions for rearrangement to form subgrains and (d) cooling conditions for suppressing coarsening of the formed fine ferrite crystal grains and fine subgrain grains.
  • Patent Document 4 discloses that the (110) plane X intensity ratio (X-ray plane intensity ratio in the (110) plane showing a texture developing degree) is 2 or more by controlled rolling and the equivalent circle diameter (diameter equivalent). To a circle in the crystal grains) It is described that the brittle fracture resistance is improved by making coarse grains of 20 ⁇ m or more 10% or less.
  • Patent Document 5 is characterized in that, as a welded structural steel having excellent brittle crack propagation stopping performance in a joint part, the (100) plane X-ray plane strength ratio in the rolled surface inside the plate thickness is 1.5 or more. Steel sheet is disclosed, and it is described that excellent brittle crack propagation stopping characteristics can be obtained by the deviation of the angle between the stress load direction and the crack propagation direction due to the texture development.
  • Japanese Patent Publication No. 7-100814 JP 2002-256375 A Japanese Patent No. 3467767 Japanese Patent No. 3548349 Japanese Patent No. 2659661 Japanese Patent No. 3546308
  • Non-Patent Document 1 evaluates the brittle crack propagation stopping performance of a steel plate having a thickness of 65 mm, and reports a result that the brittle crack does not stop in a large-scale brittle crack propagation stopping test of the base material.
  • the Kca value at the use temperature of ⁇ 10 ° C. (hereinafter also referred to as Kca ( ⁇ 10 ° C.)) satisfies 3000 N / mm 3/2 .
  • Kca ( ⁇ 10 ° C.) satisfies 3000 N / mm 3/2 .
  • the steel sheet having a thickness of about 50 mm is the main target of the steel sheets having excellent brittle crack propagation stopping characteristics described in Patent Documents 1 to 5 described above.
  • Patent Documents 1 to 5 When the techniques described in Patent Documents 1 to 5 are applied to a thick material exceeding 50 mm, it is unclear whether the predetermined characteristics can be obtained, and the characteristics against crack propagation in the plate thickness direction necessary for the hull structure are completely different. Not verified.
  • the welding work requires high efficiency such as submerged arc welding, electrogas welding, electroslag welding, etc.
  • Heat input welding is applied.
  • the structure of the weld heat-affected zone (HeatffAffected; Zone; HAZ) becomes coarse, so that the toughness of the weld heat-affected zone decreases.
  • steel materials for high heat input welding have already been developed and put to practical use.
  • Patent Document 6 by controlling TiN precipitated in steel, it prevents coarsening of the weld heat affected zone structure and promotes intragranular ferrite transformation by dispersion of ferrite forming nuclei.
  • a technique for increasing the toughness of the weld heat affected zone is disclosed.
  • the toughness of the weld heat-affected zone of the high heat input weld zone is excellent, the brittle crack propagation stop property is not taken into consideration, and those satisfying both properties have not been obtained.
  • the present invention optimizes the steel composition and rolling conditions, controls the texture in the thickness direction, and has high brittle crack propagation stopping characteristics that can be stably manufactured by an extremely simple process industrially.
  • An object of the present invention is to provide a high-strength thick steel plate and a method for producing the same.
  • FIGS. 1A and 1B are diagrams schematically showing an example in which propagation of a crack 3 entering from a notch 2 of a standard ESSO test piece 1 stops at a tip shape 4 in a base material 5, and is schematically shown in FIG. It was confirmed that high arrestability was obtained when the short crack branch 3a as shown was confirmed. It is presumed that the stress is relieved by the crack branch 3a. 2.
  • the steel structure mainly composed of bainite in which a packet or the like is present is more advantageous than the steel structure mainly composed of ferrite, and the (100) surface which is a cleavage plane is propagated by cracks. It is effective to accumulate at an angle with respect to the rolling direction or the sheet width direction. 3.
  • the degree of integration on the (100) plane is too high, a large crack branch is generated from a very short crack branch.
  • Non-Patent Document 2 showing the design guideline for brittle crack arrest of ship structure, it is necessary to suppress the branching of brittle cracks in the standard ESSO test.
  • the texture at the center of the plate thickness is controlled by performing rolling in which the difference between the rolling temperature of the first pass and the rolling temperature of the last pass is 40 ° C. or less by controlling the texture at the center of the plate thickness to 40 to 70%. Can be realized. 6).
  • the composite sulfides of TiN, CaS and MnS are finely divided to suppress grain growth when exposed to high temperatures of welding, and in the subsequent cooling process It is effective to promote internal transformation and refine the heat affected zone structure at room temperature.
  • the present invention has been made by further study based on the obtained knowledge. That is, the present invention 1.
  • Steel composition is mass%, C: 0.03-0.15%, Si: 0.01-0.5%, Mn: 1.40-2.50%, Al: 0.005-0.08 %, P: 0.03% or less, S: 0.0005 to 0.0030%, N: 0.0036 to 0.0070%, Ti: 0.004 to 0.030%, Ca: 0.0005 to 0 .0030%, and each content of Ca, S, and O satisfies the following formula (1), the balance is Fe and inevitable impurities, the metal structure is mainly bainite, and the central portion of the plate thickness
  • the RD // (110) plane has a texture of 1.5 to 4.0, and the Charpy fracture surface transition temperature vTrs in the surface layer portion and the thickness center portion is ⁇ 40 ° C.
  • High strength thick steel plate for high heat input welding with excellent brittle crack propagation stopping characteristics. 0.30 ⁇ (Ca ⁇ (0.18 + 130 ⁇ Ca) ⁇ O) /1.25/S ⁇ 0.80 (1) However, in Formula (1), Ca, O, and S are made into content (mass%). 2.
  • the steel composition is further mass%, Nb: 0.05% or less, Cu: 1.0% or less, Ni: 1.0% or less, Cr: 0.5% or less, Mo: 0.5% or less, V: 0.2% or less, B: 0.003% or less, REM: 0.01% or less of one type or two or more types, which is excellent in brittle crack propagation stop property according to 1. High strength thick steel plate for high heat input welding. 3.
  • t is a plate thickness (mm).
  • the steel material having the composition described in 4.1 or 2 is heated to a temperature of 1000 to 1200 ° C., and rolling with a total cumulative rolling reduction of 65% or more in the austenite recrystallization temperature range and the austenite non-recrystallization temperature range is performed.
  • the rolling reduction is performed at a cumulative reduction ratio of 20% or more, and then in the state where the plate thickness central portion is in the austenite non-recrystallization temperature range.
  • the heat input is excellent in brittle crack propagation stopping characteristics, characterized in that the difference is within 40 ° C. and then cooled to 450 ° C. or lower at a cooling rate of 4.0 ° C./s or higher.
  • Method of producing a high strength steel plate for contact. 5 After accelerated cooling to 450 ° C. or lower, and further tempering to a temperature of A c1 point or lower, the high strength thickness for high heat input welding having excellent brittle crack propagation stopping characteristics according to 4 A method of manufacturing a steel sheet.
  • the present invention it is possible to obtain a high-strength thick steel plate and a method for producing the same, in which the texture is appropriately controlled in the plate thickness direction, and the brittle crack propagation stop property and the high heat input weld joint toughness are excellent.
  • Applying the present invention to a steel plate having a plate thickness of 50 mm or more, preferably more than 50 mm, more preferably 55 mm or more, and even more preferably 60 mm or more is more significant than the steel according to the prior art. It is effective because it demonstrates its properties. And, for example, in the shipbuilding field, it contributes to improving the safety of ships by applying the present invention to hatch side combing and deck members in the structure of large container ships and bulk carrier strong deck parts. .
  • FIG. 1 is a diagram schematically showing a fracture surface form of a standard ESSO test of a thick steel plate having a thickness of more than 50 mm, (a) is a view of the test piece observed from the plane side, and (b) is a fracture of the test piece. It is a figure which shows a surface.
  • C 0.03-0.15%
  • C is an element that improves the strength of steel, and in the present invention, it is necessary to contain 0.03% or more in order to ensure a desired strength. On the other hand, if it exceeds 0.15%, the weldability is deteriorated and the toughness is also adversely affected. Therefore, C is specified in the range of 0.03 to 0.15%. Preferably, it is 0.05 to 0.15%.
  • Si 0.01 to 0.5% Si is effective as a deoxidizing element and as a steel strengthening element. However, when the content is less than 0.01%, the effect is not obtained. On the other hand, if it exceeds 0.5%, not only the surface properties of the steel are impaired, but also the toughness is extremely deteriorated. Therefore, the addition amount is set to 0.01 to 0.5%. Preferably, it is 0.02 to 0.45% of range.
  • Mn 1.40-2.50% Mn is added as a strengthening element. If less than 1.40%, the effect is not sufficient. On the other hand, if it exceeds 2.50%, the weldability deteriorates and the steel material cost also increases. Therefore, Mn is set to 1.40 to 2.50%. Preferably, it is in the range of 1.42 to 2.40%.
  • P 0.03% or less
  • the toughness of the welded portion is significantly deteriorated.
  • the upper limit is made 0.03%.
  • it is 0.02% or less.
  • S 0.0005 to 0.0030%
  • S is required to be 0.0005% or more in order to generate necessary CaS and MnS.
  • S is set to 0.0005 to 0.0030%.
  • it is in the range of 0.0006 to 0.0025%.
  • Al acts as a deoxidizing agent, and for this purpose, a content of 0.005% or more is required. However, when it contains exceeding 0.08%, while reducing toughness, when welding, the toughness of a weld metal part will be reduced. For this reason, Al is specified in the range of 0.005 to 0.08%. Preferably, it is 0.02 to 0.06%.
  • Ti can form nitrides, carbides, or carbonitrides by adding a small amount, suppress austenite coarsening in the weld heat affected zone, and / or promote ferrite transformation as a ferrite transformation nucleus. This has the effect of refining the crystal grains and improving the base material toughness. The effect is obtained by adding 0.004% or more. However, the content exceeding 0.030% reduces the toughness of the base material and the weld heat affected zone due to the coarsening of the TiN particles. Therefore, Ti is set in the range of 0.004 to 0.030%. Preferably, it is 0.006 to 0.028% of range.
  • N 0.0036 to 0.0070%
  • N is an element necessary for securing the necessary amount of TiN. If it is less than 0.0036%, a sufficient amount of TiN cannot be obtained, and the weld toughness deteriorates. If it exceeds 0.0070%, TiN will re-dissolve when subjected to the welding heat cycle, and excessive N will be generated and the toughness will deteriorate significantly. For this reason, N is made 0.0036 to 0.0070%. Preferably, it is 0.0038 to 0.0065% of range.
  • Ca 0.0005 to 0.0030%
  • Ca is an element having an effect of improving toughness by fixing S. In order to exhibit such an effect, it is necessary to contain at least 0.0005% or more. However, the effect is saturated even if the content exceeds 0.0030%. Therefore, in the present invention, Ca is limited to the range of 0.0005 to 0.0030%. Preferably, it is in the range of 0.0007 to 0.0028%.
  • the above is the basic component composition of the present invention.
  • Nb 0.05% or less Nb precipitates as NbC during ferrite transformation or reheating, and contributes to increasing the strength. In addition, it has the effect of expanding the non-recrystallization temperature range in rolling in the austenite range, and contributes to the fine graining of bainite packets, so it is also effective in improving toughness. Since the effect is exhibited by containing 0.005% or more, when it contains, it is preferable to make it 0.005% or more. However, if added over 0.05%, coarse NbC precipitates and conversely causes a decrease in toughness. When it is contained, the upper limit is preferably made 0.05%. More preferably, it is in the range of 0.007 to 0.045%.
  • Cu, Ni, Cr, Mo Cu, Ni, Cr, and Mo are all elements that enhance the hardenability of steel. While contributing directly to strength enhancement after rolling, it can be added to improve functions such as toughness, high-temperature strength, or weather resistance, since these effects are exhibited by containing 0.01% or more, When contained, the content is preferably 0.01% or more. However, if it is excessively contained, toughness and weldability are deteriorated. Therefore, when it is included, the upper limit is 1.0% for Cu, 1.0% for Ni, 0.5% for Cr, and 0.5% for Mo. % Is preferable. More preferably, Cu: 0.02 to 0.95%, Ni: 0.02 to 0.95%, Cr: 0.02 to 0.46%, Mo: 0.02 to 0.46% is there.
  • V 0.2% or less
  • V is an element that improves the strength of steel by precipitation strengthening as V (C, N), and may be contained by 0.001% or more in order to exert this effect. However, when it contains exceeding 0.2%, toughness will be reduced. Therefore, when V is contained, the content is preferably 0.2% or less, and more preferably in the range of 0.001 to 0.10%.
  • B 0.003% or less
  • B is an element that enhances the hardenability of steel in a small amount, and may be contained by 0.0005% or more in order to exert this effect. However, if it exceeds 0.003%, the toughness of the welded portion is lowered. Therefore, when B is contained, the content is preferably 0.003% or less. More preferably, it is in the range of 0.0006 to 0.0025%.
  • REM 0.01% or less REM refines the structure of the weld heat-affected zone to improve toughness, and even if added, the effect of the present invention is not impaired, so it may be added as necessary. Since this effect is exhibited by containing 0.0010% or more, when it is contained, the content is preferably 0.0010% or more. However, if added excessively, coarse inclusions are formed and the toughness of the base material is deteriorated. Therefore, when added, the upper limit of the addition amount is preferably 0.01%.
  • O is contained in steel as an unavoidable impurity and reduces cleanliness. For this reason, in the present invention, it is desirable to reduce O as much as possible.
  • the O content exceeds 0.0050%, CaO inclusions are coarsened and the base material toughness is lowered. For this reason, Preferably it is 0.0050% or less.
  • the present invention in order to crystallize Ca as CaS, it is necessary to reduce the amount of O having strong binding force with Ca before adding Ca, and the residual oxygen amount before adding Ca is 0.0050%.
  • the following is preferable.
  • a method for reducing the amount of residual oxygen a method such as enhancing degassing or introducing a deoxidizer can be employed.
  • the balance other than the above components is Fe and inevitable impurities.
  • the brittle crack propagation stop property is exhibited for cracks that develop in the horizontal direction (in-plane direction of the steel sheet) such as the rolling direction or the direction perpendicular to the rolling direction.
  • the toughness at the surface thickness layer and the central portion and the degree of integration of the RD // (100) plane at the central portion of the thickness are appropriately defined according to the desired brittle crack propagation stop characteristics.
  • the Charpy fracture surface transition temperature at the plate thickness surface layer portion and the central portion is defined as ⁇ 40 ° C. or less as the toughness at the plate thickness surface layer portion and the center portion.
  • the Charpy fracture surface transition temperature at the center of the plate thickness is preferably ⁇ 50 ° C. or lower.
  • the cleavage plane is accumulated obliquely with respect to the main crack direction, and the effect of stress relaxation at the brittle crack tip by generating fine crack branching causes brittleness.
  • the crack propagation stop performance is improved.
  • the degree of integration of the RD // (110) plane in the central portion of the plate thickness needs to be 1.5 or more, preferably 1.7 or more. Therefore, in the present invention, the degree of integration of the RD // (110) plane at the center of the plate thickness is 1.5 or more, preferably 1.7 or more.
  • the integration degree of the RD // (110) plane is set to a range of 1.5 to 4.0.
  • the degree of integration of the RD // (110) plane in the central portion of the plate thickness refers to the following.
  • a sample with a plate thickness of 1 mm is taken from the center of the plate thickness, and a test piece for X-ray diffraction is prepared by mechanically polishing and electrolytic polishing a surface parallel to the plate surface.
  • an X-ray diffraction measurement was performed using an X-ray diffractometer using a Mo ray source, and (200), (110) and (211) positive electrode dot diagrams were obtained and obtained.
  • a three-dimensional crystal orientation density function is calculated from the positive electrode dot diagram by the Bunge method.
  • the integrated value is obtained by integrating the values of the three-dimensional crystal orientation density function of the orientation, and the value obtained by dividing the integrated value by the number of the integrated orientations is referred to as the degree of integration of the RD // (110) plane.
  • the Charpy fracture surface transition temperature at the center of the plate thickness and the degree of integration of the RD // (110) plane satisfy the following formula (2) in addition to the above-mentioned provisions of the base material toughness and texture.
  • formula (2) further excellent brittle crack propagation stopping performance can be obtained.
  • vTrs (1 / 2t) Charpy fracture surface transition temperature (° C.) at the center of the plate thickness I RD // (110) [1 / 2t] : RD // (110) integration degree at the center of the plate thickness.
  • t is a plate thickness (mm).
  • the degree of integration of the RD // (110) plane can be 1.5 or more, preferably 1.7 or more.
  • the metal structure obtained after rolling and cooling is mainly bainite.
  • that the metal structure is mainly bainite is that the area fraction of the bainite phase is 80% or more of the whole. The balance is acceptable if ferrite, martensite (including island martensite), pearlite, etc. are 20% or less in total area fraction.
  • the heating temperature, hot rolling conditions, cooling conditions, etc. of the steel material As manufacturing conditions, it is preferable to prescribe the heating temperature, hot rolling conditions, cooling conditions, etc. of the steel material.
  • hot rolling in addition to the cumulative reduction ratio in the sum of the austenite recrystallization temperature range and the austenite non-recrystallization temperature range, the case where the central portion of the plate thickness is in the austenite recrystallization temperature range, It is preferable to define the cumulative rolling reduction for each of the cases in the temperature range and the rolling temperature conditions in a state where the central portion of the plate thickness is in the austenite non-recrystallized region.
  • the Charpy fracture surface transition temperature vTrs in the surface layer portion and the plate thickness center portion of the thick steel plate, and the RD // (110) integration degree in the plate thickness center portion can be set to desired values.
  • molten steel having the above composition is melted in a converter or the like, and is made into a steel material (slab) by continuous casting or the like.
  • the heating temperature is preferably 1000 to 1200 ° C.
  • a more preferable heating temperature range is 1000 to 1150 ° C. from the viewpoint of toughness.
  • the degree of integration of the RD // (110) plane can be 1.5 or more, preferably 1.7 or more.
  • the hot rolling first, it is preferable to perform rolling with a cumulative reduction ratio of 20% or more in a state where the central portion of the plate thickness is in the austenite recrystallization temperature region.
  • a cumulative reduction ratio of 20% or more in a state where the central portion of the plate thickness is in the austenite recrystallization temperature region.
  • the cumulative reduction ratio of 40 to 70% or more in a state where the temperature at the center of the plate thickness is in the austenite non-recrystallization temperature range.
  • the cumulative reduction ratio in this temperature range is 40% or more, the texture at the center of the plate thickness is sufficiently developed, and the degree of integration of the RD // (110) plane at the center of the plate thickness is 1.5 or more. , And preferably 1.7 or more.
  • the range of the cumulative rolling reduction is set to 40 to 70%.
  • the rolling temperature refers to the temperature at the center of the plate thickness of the steel just before rolling.
  • the temperature at the center of the plate thickness is obtained by simulation calculation or the like from the plate thickness, surface temperature, thermal history, and the like. For example, the temperature at the center of the plate thickness of the steel sheet is obtained by calculating the temperature distribution in the plate thickness direction using the difference method.
  • the total cumulative rolling reduction of the austenite recrystallization temperature range and the austenite non-recrystallization temperature range be 65% or more.
  • the overall rolling reduction is small, the rolling of the structure is not sufficient, and the toughness and strength cannot achieve the target values.
  • the total cumulative reduction ratio By setting the total cumulative reduction ratio to 65% or more, a sufficient amount of reduction can be ensured for the structure, and the toughness and the degree of accumulation can achieve the target values.
  • the austenite recrystallization temperature range and the austenite non-recrystallization temperature range can be grasped by conducting a preliminary experiment in which the steel having the component composition is given a heat / working history with varying conditions.
  • end temperature of hot rolling is not particularly limited. From the viewpoint of rolling efficiency, it is preferable to terminate in the austenite non-recrystallization temperature range.
  • the rolled steel sheet is cooled to 450 ° C. or lower at a cooling rate of 4.0 ° C./s or higher.
  • the cooling rate is less than 4.0 ° C./s, the coarsening of the structure and ferrite transformation proceed at each plate thickness position, so that a desired structure cannot be obtained and the strength of the steel sheet also decreases.
  • the bainite transformation can be sufficiently advanced, and desired toughness and integration degree can be obtained. If the cooling stop temperature is higher than 450 ° C., the bainite transformation does not proceed sufficiently, and a structure such as ferrite or pearlite is also produced, and the bainite-based structure intended by the present invention cannot be obtained.
  • these cooling rate and cooling stop temperature be the temperature of the plate
  • Tempering temperature as follows C1 points A steel plate average temperature, by carrying out the tempering treatment, it is possible not impair the desired tissue obtained by rolling and cooling.
  • the AC1 point (° C.) is obtained by the following equation.
  • a C1 point 751-26.6C + 17.6Si-11.6Mn-169Al-23Cu-23Ni + 24.1Cr + 22.5Mo + 233Nb-39.7V-5.7Ti-895B
  • each element symbol is the content (% by mass) in steel, and 0 if not contained.
  • the average temperature of the steel sheet can also be obtained by simulation calculation or the like from the sheet thickness, surface temperature, cooling conditions, etc., similarly to the temperature at the center of the sheet thickness.
  • Molten steel (steel symbols A to Q) of each composition shown in Table 1 was melted in a converter and made into a steel material (slab 250 mm thickness or 300 mm thickness) by a continuous casting method. After hot rolling to a plate thickness of 55 to 100 mm, Cooling is performed, and no. Sample steels of 1 to 27 were obtained. Some were tempered after cooling. Table 2 shows hot rolling conditions and cooling conditions.
  • a JIS14A test piece having a diameter of 14 mm was collected from 1/4 part of the plate thickness so that the longitudinal direction of the test piece was perpendicular to the rolling direction, a tensile test was performed, and the yield strength (YS) and Tensile strength (TS) was measured.
  • the longitudinal direction of the test piece was measured in accordance with the JIS No. 4 impact test piece from the plate thickness surface layer portion and the plate thickness central portion (hereinafter, the plate thickness central portion may be referred to as a 1/2 t portion). Samples were taken so as to be parallel to the rolling direction, and Charpy impact tests were performed to determine fracture surface transition temperatures (vTrs).
  • the impact test piece of the surface layer part is assumed to have a surface closest to the surface at a depth of 1 mm from the steel sheet surface.
  • the degree of integration of the RD // (110) plane at the center of the plate thickness was determined as follows. First, a sample having a plate thickness of 1 mm was collected from the central portion of the plate thickness, and a test piece for X-ray diffraction was prepared by mechanically polishing and electrolytic polishing a surface parallel to the plate surface. Using this test piece, X-ray diffraction measurement was performed using an X-ray diffractometer using a Mo ray source, and (200), (110) and (211) positive electrode dot diagrams were obtained, and the obtained positive electrode A three-dimensional crystal orientation distribution density function is calculated from the dot diagram by the Bunge method.
  • test steel sheet was subjected to groove processing (groove angle 20 °), and heat input 300 was performed by electrogas welding using a commercially available wire for electrogas arc welding for low temperature steel.
  • a welded joint was prepared at ⁇ 750 kJ / cm, and the toughness of the bond portion was evaluated by a 2 mmV notch Charpy test as HAZ toughness.
  • the test was performed with vE -20 (average value of three) of Charpy absorbed energy at -20 ° C.
  • Table 3 shows the results of these tests.
  • the test steel plates (production Nos. 1 to 11) within the scope of the present invention exhibited excellent brittle crack propagation stopping performance with a Kca ( ⁇ 10 ° C.) of 6000 N / mm 3/2 or more. Further, the absorbed energy of the bond portion of the high heat input welded joint was vE-20 ⁇ 88 J, which was an excellent value. Further, in the test steel plates (manufacturing numbers 2 to 11) in which the Charpy toughness value (fracture surface transition temperature) of the surface layer portion and the center portion of the plate thickness and the RD // (110) accumulation degree satisfy the formula (2) Compared with the test steel plate (Production No. 1) not satisfying the formula (2), a higher Kca ( ⁇ 10 ° C.) value was obtained. Note that the metal structures of these test steel sheets (production Nos. 1 to 11) were mainly bainite.
  • the steel plate component is within the preferred range of the present invention
  • the steel plate (production Nos. 20 to 27) whose heating and rolling conditions are outside the preferred range of the present invention is Kca ( ⁇ 10 ° C.). The value did not reach 6000 N / mm 3/2 .
  • the absorbed energy of the large heat input welded joint: vE- 20 is 22 J or less, which is inferior to the present invention example. It was.

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Abstract

Provided are a high strength thick steel plate for high heat input welding, the plate of a preferred plate thickness of at least 50 mm being used for ships and having excellent brittle crack arrestability, and a manufacturing method therefor. A thick steel plate having a specific component composition, the main constituent of the metal structure being bainite, the thick steel plate having a texture in which the density in the RD//(110) plane at the plate thickness center is 1.5-4.0, and the Charpy fracture appearance transition temperature (vTrs) at the surface layer and the plate thickness center being no more than -40°C; and a manufacturing method therefor.

Description

脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板およびその製造方法High strength thick steel plate for high heat input welding with excellent brittle crack propagation stopping characteristics and method for producing the same
　本発明は、脆性亀裂伝播停止特性(brittle crack arrestability)に優れた大入熱溶接(high heat input welding)用高強度厚鋼板(high-strength thick steel plate)およびその製造方法に関し、特に、船舶に用いて好適な板厚５０ｍｍ以上のものに関する。 The present invention relates to a high-strength thick steel plate for high-heat input welding having excellent brittle cracking arrestability and a method for producing the same, and more particularly to a ship. It relates to a plate having a thickness of 50 mm or more suitable for use.
　船舶等の大型構造物においては、脆性破壊(brittle fracture)に伴う事故が経済や環境に及ぼす影響が大きい。このため、安全性の向上が常に求められ、使用される鋼材に対しては、使用温度における靭性(toughness)や、脆性亀裂伝播停止特性が要求されている。 In large structures such as ships, accidents associated with brittle fractures have a great impact on the economy and the environment. For this reason, improvement of safety is always required, and steel materials to be used are required toughness at use temperature and brittle crack propagation stop characteristics.
　コンテナ船やバルクキャリアーなどの船舶はその構造上、船体外板(outer plate of ship's hull)に高強度の厚肉材を使用する。最近は船体の大型化に伴い一層の高強度厚肉化が進展し、一般に、鋼板の脆性亀裂伝播停止特性は高強度あるいは厚肉材ほど劣化する傾向があるため、脆性亀裂伝播停止特性への要求も一段と高度化している。 Ships such as container ships and bulk carriers use high-strength, thick materials for the outer shell plate (outer plate plate of ship's hull). Recently, as the hull size has increased, the strength and thickness of steel sheets have further increased, and generally, the brittle crack propagation stopping characteristics of steel sheets tend to deteriorate as the strength and thickness of the steel increases. Requests are becoming more sophisticated.
　鋼材の脆性亀裂伝播停止特性を向上させる手段として、従来からＮｉ含有量を増加させる方法が知られており、液化天然ガス（ＬＮＧ：Liquefied Natural Gas）の貯槽タンクにおいては、９％Ｎｉ鋼が商業規模で使用されている。 As a means of improving the brittle crack propagation stopping characteristics of steel materials, a method of increasing the Ni content has been conventionally known. In a liquefied natural gas (LNG) storage tank, 9% Ni steel is commercially available. Used on a scale.
　しかし、Ｎｉ量の増加はコストの大幅な上昇を余儀なくさせるため、ＬＮＧ貯槽タンク以外の用途には適用が難しい。 However, an increase in the amount of Ni necessitates a significant increase in cost, making it difficult to apply to applications other than LNG storage tanks.
　一方、ＬＮＧのような極低温(ultra low temperature)にまで至らない、船舶やラインパイプに使用される、板厚が５０ｍｍ未満の比較的薄手の鋼材に対しては、ＴＭＣＰ(Thermo-Mechanical Control Process)法により細粒化を図り、低温靭性を向上させることにより、優れた脆性亀裂伝播停止特性を付与することができる。 On the other hand, TMCP (Thermo-Mechanical Control Process) is used for relatively thin steel materials with a thickness of less than 50 mm used for ships and line pipes that do not reach ultra low temperatures such as LNG. ) Method to improve graininess and improve low temperature toughness, it is possible to impart excellent brittle crack propagation stopping characteristics.
　また、合金コストを上昇させることなく、脆性亀裂伝播停止特性を向上させるために、表層部の組織を超微細化(ultra fine grained steel)した鋼材が特許文献１で提案されている。 Further, in order to improve the brittle crack propagation stop characteristic without increasing the alloy cost, a steel material in which the structure of the surface layer portion is ultrafine (ultra-fine-grained-steel) is proposed in Patent Document 1.
　特許文献１には、脆性亀裂が伝播する際、鋼材表層部に発生するシアリップ（塑性変形領 shear-lips）が脆性亀裂伝播停止特性の向上に効果があることに着目し、シアリップ部分の結晶粒を微細化させて、伝播する脆性亀裂が有する伝播エネルギーを吸収させることを特徴とする脆性亀裂伝播停止特性に優れた鋼材が記載されている。 Patent Document 1 focuses on the fact that shear lips (plastic deformation regions shear-lips) that occur in the steel surface layer when brittle cracks propagate are effective in improving the brittle crack propagation stopping characteristics. A steel material excellent in brittle crack propagation stopping characteristics characterized in that the propagation energy possessed by the propagating brittle cracks is absorbed is described.
　また、特許文献１には、熱間圧延後の制御冷却により表層部分をＡｒ３変態点(transformation point)以下に冷却し、その後制御冷却(controlled cooling)を停止して表層部分を変態点以上に復熱(recuperate)させる工程を１回以上繰り返して行い、この間に鋼材に圧下を加えることにより、繰り返し変態させ又は加工再結晶させて、表層部分に超微細なフェライト組織(ferrite structure)又はベイナイト組織(bainite structure)を生成させることが記載されている。 Further, Patent Document 1 discloses that the surface layer portion is cooled below the Ar3 transformation point by controlled cooling after hot rolling, and then the controlled cooling is stopped to bring the surface layer portion above the transformation point. The process of recuperate is repeated one or more times, and during this time, the steel material is subjected to reduction, and it is repeatedly transformed or processed and recrystallized, so that a superfine ferrite structure or bainite structure is formed on the surface layer portion. (bainite structure) is generated.
　さらに、特許文献２では、フェライト－パーライト(pearlite)を主体のミクロ組織とする鋼材において脆性亀裂伝播停止特性を向上させるためには、鋼材の両表面部は円相当粒径(circle-equivalent average grain size)：５μｍ以下、アスペクト比(aspect ratio of the grains)：２以上のフェライト粒を有するフェライト組織を５０％以上有する層で構成し、フェライト粒径のバラツキを抑えることが重要であるとしている。バラツキを抑える方法として仕上げ圧延中の１パス当りの最大圧下率(maximum rolling reduction)を１２％以下とし局所的な再結晶現象を抑制することが記載されている。 Furthermore, in Patent Document 2, in order to improve the brittle crack propagation stop property in a steel material mainly composed of ferrite-pearlite, both surface portions of the steel material have a circle-equivalent average grain. size): 5 μm or less, aspect ratio (of aspect ratio): a layer having 50% or more of a ferrite structure having two or more ferrite grains, and suppressing variation in ferrite grain size is important. As a method for suppressing the variation, it is described that the maximum rolling reduction per pass during finish rolling is set to 12% or less to suppress the local recrystallization phenomenon.
　しかし、特許文献１、２に記載の脆性亀裂伝播停止特性に優れた鋼材は、鋼材表層部のみを一旦冷却した後に復熱させ、かつ復熱中に加工を加えることによって、特定の組織を得るものであるため、実生産規模では制御が容易ではない。特に板厚が５０ｍｍを超える厚肉材では、圧延、冷却設備への負荷が大きいプロセスである。 However, the steel materials excellent in brittle crack propagation stopping characteristics described in Patent Documents 1 and 2 are obtained by recooling only the steel surface layer part and then recovering the heat, and by applying processing during the recuperation, a specific structure is obtained. Therefore, control is not easy on the actual production scale. In particular, a thick material having a plate thickness exceeding 50 mm is a process with a heavy load on rolling and cooling equipment.
　一方、特許文献３には、フェライト結晶粒の微細化のみならずフェライト結晶粒内に形成されるサブグレイン(subgrain)に着目し、脆性亀裂伝播停止特性を向上させる、ＴＭＣＰの延長上にある技術が記載されている。 On the other hand, in Patent Document 3, attention is paid not only to the refinement of ferrite crystal grains but also to subgrains formed in ferrite crystal grains, and a technique on the extension of TMCP that improves brittle crack propagation stop characteristics. Is described.
　具体的には、板厚３０～４０ｍｍの鋼板において、鋼板表層の冷却および復熱などの複雑な温度制御を必要とせずに、（ａ）微細なフェライト結晶粒を確保する圧延条件、（ｂ）鋼材板厚の５％以上の部分に微細フェライト組織を生成する圧延条件、（ｃ）微細フェライトに集合組織(texture)を発達させるとともに加工（圧延）により導入した転位(dislocation)を熱的エネルギーにより再配置しサブグレインを形成させる圧延条件、（ｄ）形成した微細なフェライト結晶粒と微細なサブグレイン粒の粗大化を抑制する冷却条件、によって脆性亀裂伝播停止特性を向上させる。 Specifically, in a steel sheet having a thickness of 30 to 40 mm, (a) rolling conditions for securing fine ferrite crystal grains without requiring complicated temperature control such as cooling and recuperation of the steel sheet surface layer, (b) Rolling conditions for generating a fine ferrite structure in a portion of 5% or more of the steel sheet thickness, (c) Dislocation introduced by machining (rolling) and development of texture in the fine ferrite by thermal energy The brittle crack propagation stop property is improved by rolling conditions for rearrangement to form subgrains and (d) cooling conditions for suppressing coarsening of the formed fine ferrite crystal grains and fine subgrain grains.
　また、制御圧延において、変態したフェライトに圧下を加えて集合組織を発達させることにより、脆性亀裂伝播停止特性を向上させる方法も知られている。この方法は、鋼材の破壊面上にセパレーション(separation)を板面と平行な方向に生ぜしめ、脆性亀裂先端の応力を緩和させることにより、脆性破壊に対する抵抗を高める。 Also known is a method of improving the brittle crack propagation stop property by controlling the rolling to develop a texture by reducing the transformed ferrite. This method increases the resistance to brittle fracture by causing separation on the fracture surface of the steel material in a direction parallel to the plate surface and relaxing stress at the tip of the brittle crack.
　例えば、特許文献４には、制御圧延により（１１０）面Ｘ線強度比(X-ray plane intensity ratio in the (110) plane showing a texture developing degree)を２以上とし、かつ円相当径(diameter equivalent to a circle in the crystal grains)２０μｍ以上の粗大粒を１０％以下とすることにより、耐脆性破壊特性を向上させることが記載されている。 For example, Patent Document 4 discloses that the (110) plane X intensity ratio (X-ray plane intensity ratio in the (110) plane showing a texture developing degree) is 2 or more by controlled rolling and the equivalent circle diameter (diameter equivalent). To a circle in the crystal grains) It is described that the brittle fracture resistance is improved by making coarse grains of 20 μm or more 10% or less.
　特許文献５には継手部の脆性亀裂伝播停止性能の優れた溶接構造用鋼として、板厚内部の圧延面における（１００）面のＸ線面強度比が１．５以上を有することを特徴とする鋼板が開示され、当該集合組織発達による応力負荷方向と亀裂伝播方向の角度のずれにより優れた脆性亀裂伝播停止特性が得られることが記載されている。 Patent Document 5 is characterized in that, as a welded structural steel having excellent brittle crack propagation stopping performance in a joint part, the (100) plane X-ray plane strength ratio in the rolled surface inside the plate thickness is 1.5 or more. Steel sheet is disclosed, and it is described that excellent brittle crack propagation stopping characteristics can be obtained by the deviation of the angle between the stress load direction and the crack propagation direction due to the texture development.
特公平７－１００８１４号公報Japanese Patent Publication No. 7-100814 特開２００２－２５６３７５号公報JP 2002-256375 A 特許第３４６７７６７号公報Japanese Patent No. 3467767 特許第３５４８３４９号公報Japanese Patent No. 3548349 特許第２６５９６６１号公報Japanese Patent No. 2659661 特許第３５４６３０８号公報Japanese Patent No. 3546308
　ところで、最近の６、０００ＴＥＵ(Twenty-foot Equivalent Unit)を超える大型コンテナ船では板厚５０ｍｍを超える厚鋼板が使用される。非特許文献１は、板厚６５ｍｍの鋼板の脆性亀裂伝播停止性能を評価し、母材の大型脆性亀裂伝播停止試験で脆性亀裂が停止しない結果を報告している。 By the way, in the recent large container ships exceeding 6,000 TEU (Twenty-foot Equivalent Unit), thick steel plates exceeding 50 mm are used. Non-Patent Document 1 evaluates the brittle crack propagation stopping performance of a steel plate having a thickness of 65 mm, and reports a result that the brittle crack does not stop in a large-scale brittle crack propagation stopping test of the base material.
　また、供試材の標準ＥＳＳＯ試験(ESSO test compliant with WES 3003)では使用温度－１０℃におけるＫｃａの値（以下、Ｋｃａ（－１０℃）とも記載する。）が３０００Ｎ／ｍｍ３／２に満たない結果が示され、５０ｍｍを超える板厚の鋼板を適用した船体構造の場合、安全性確保が課題となることが示唆されている。 Further, in the standard ESSO test (ESSO test compliant with WES 3003) of the test material, the Kca value at the use temperature of −10 ° C. (hereinafter also referred to as Kca (−10 ° C.)) satisfies 3000 N / mm 3/2 . In the case of a hull structure to which a steel plate having a thickness exceeding 50 mm is applied, it is suggested that ensuring safety is an issue.
　上述した特許文献１～５に記載の脆性亀裂伝播停止特性に優れる鋼板は、製造条件や開示されている実験データから、板厚５０ｍｍ程度までの鋼板が主な対象であると考えられる。特許文献１～５記載の技術を５０ｍｍを超える厚肉材へ適用した場合、所定の特性が得られるか不明であり、船体構造で必要な板厚方向の亀裂伝播に対しての特性については全く検証されていない。 From the manufacturing conditions and the disclosed experimental data, it is considered that the steel sheet having a thickness of about 50 mm is the main target of the steel sheets having excellent brittle crack propagation stopping characteristics described in Patent Documents 1 to 5 described above. When the techniques described in Patent Documents 1 to 5 are applied to a thick material exceeding 50 mm, it is unclear whether the predetermined characteristics can be obtained, and the characteristics against crack propagation in the plate thickness direction necessary for the hull structure are completely different. Not verified.
　一方、鋼板の厚肉化にともない、溶接施工には、サブマージアーク溶接(submerged arc welding)、エレクトロガス溶接(electrogas arc welding)、エレクトロスラグ溶接(electroslag welding)などの高能率(high efficiency)な大入熱溶接が適用されている。一般に、溶接入熱量が大きくなると、溶接熱影響部(Heat Affected Zone;HAZ)の組織が粗大化するために、溶接熱影響部の靭性は低下することが知られている。このような大入熱溶接による靭性の低下問題を解決するために、大入熱溶接用鋼材が既に開発され、実用化に至っている。例えば、特許文献６においては、鋼中に析出するＴｉＮを制御することにより溶接熱影響部組織の粗大化(coarsening)を防止するとともに、フェライト生成核の分散によって粒内フェライト変態を促進することにより、溶接熱影響部を高靭化する技術が開示されている。しかし、大入熱溶接部の溶接熱影響部の靭性には優れるものの、脆性亀裂伝播停止特性は考慮されておらず、両特性を満足するものは得られていなかった。 On the other hand, as the thickness of the steel plate increases, the welding work requires high efficiency such as submerged arc welding, electrogas welding, electroslag welding, etc. Heat input welding is applied. In general, it is known that when the heat input of welding increases, the structure of the weld heat-affected zone (HeatffAffected; Zone; HAZ) becomes coarse, so that the toughness of the weld heat-affected zone decreases. In order to solve such a problem of lowering toughness due to high heat input welding, steel materials for high heat input welding have already been developed and put to practical use. For example, in Patent Document 6, by controlling TiN precipitated in steel, it prevents coarsening of the weld heat affected zone structure and promotes intragranular ferrite transformation by dispersion of ferrite forming nuclei. A technique for increasing the toughness of the weld heat affected zone is disclosed. However, although the toughness of the weld heat-affected zone of the high heat input weld zone is excellent, the brittle crack propagation stop property is not taken into consideration, and those satisfying both properties have not been obtained.
　そこで本発明は、鋼成分、圧延条件を最適化し、板厚方向での集合組織を制御する、工業的に極めて簡易なプロセスで安定して製造し得る脆性亀裂伝播停止特性に優れる大入熱溶接用高強度厚鋼板およびその製造方法を提供することを目的とする。 Therefore, the present invention optimizes the steel composition and rolling conditions, controls the texture in the thickness direction, and has high brittle crack propagation stopping characteristics that can be stably manufactured by an extremely simple process industrially. An object of the present invention is to provide a high-strength thick steel plate and a method for producing the same.
　本発明者らは、上記課題の達成に向けて鋭意研究を重ね、厚肉鋼板でも優れた脆性亀裂伝播停止特性を有する高強度厚鋼板について以下の知見を得た。
１．板厚５０ｍｍを超える厚鋼板について、標準ＥＳＳＯ試験を行った。図１（ａ）（ｂ）は標準ＥＳＳＯ試験片１のノッチ２から突入した亀裂３が母材５において先端形状４で伝播を停止した例を模式的に示す図で（ａ）に模式的に示すような、短い亀裂の分岐３ａが確認された場合に、高いアレスト性が得られることを確認した。亀裂の分岐３ａにより応力が緩和さるためと推測される。
２．上記の破面形態を得るためには、亀裂を分岐させる組織形態にする必要がある。ここで、フェライトを主体とする鋼組織よりも、内部にパケット(packet)等が存在するベイナイトを主体とする鋼組織のほうが有利であり、また、へき開面である（１００）面を亀裂の進展方向である圧延方向あるいは板幅方向に対して斜めに集積させることが有効である。
３．一方、（１００）面の集積度を高めすぎると、極短い亀裂の分岐から、大きな亀裂の分岐が発生する。船体構造の脆性亀裂アレスト設計指針を示した非特許文献２に記載されているように、標準ＥＳＳＯ試験においては、脆性亀裂の分岐を抑制する必要があるため、亀裂の明瞭な分岐を防止するために面集積度の上限を規定する必要がある。
４．標準ＥＳＳＯ試験の破面を詳細に観察・解析した結果、亀裂の先端部となる板厚中央部の材質を制御することがアレスト性能改善に効果的であり、特に板厚中央部の靭性および集合組織に関する指標として下記（２）式をみたすことが有効である。
ｖＴｒｓ（１／２ｔ）－１２×ＩＲＤ／／（１１０）［１／２ｔ］≦－７０・・・（２）
上記式（２）において、
ｖＴｒｓ（１／２ｔ）：板厚中央部（＝１／２ｔ）の破面遷移温度（℃）
ＩＲＤ／／（１１０）［１／２ｔ］：板厚中央部（＝１／２ｔ）のＲＤ／／（１１０）面の集積度
とし、ｔは板厚（ｍｍ）である。
５．さらに、オーステナイト再結晶温度域にある状態において累積圧下率を２０％以上とする圧延を実施することによって組織の細粒化を図り、その後、オーステナイト未再結晶温度域にある状態において累積圧下率を４０～７０％とし、かつ、最初のパスの圧延温度と最後のパスの圧延温度との差が４０℃以内である圧延を実施することによって、板厚中央部の集合組織を制御して、上述の組織を実現できる。
６．大入熱溶接部の靭性を向上する手法として、ＴｉＮ，ＣａＳとＭｎＳの複合硫化物を微細に***させ、溶接の高温に曝された際の粒成長を抑制、且つ、その後の冷却過程で粒内変態を促進して室温での熱影響部組織を微細化することが有効である。 The inventors of the present invention have made extensive studies to achieve the above-mentioned problems, and have obtained the following knowledge regarding a high-strength thick steel plate having excellent brittle crack propagation stopping characteristics even with a thick steel plate.
1. A standard ESSO test was performed on a thick steel plate having a thickness exceeding 50 mm. FIGS. 1A and 1B are diagrams schematically showing an example in which propagation of a crack 3 entering from a notch 2 of a standard ESSO test piece 1 stops at a tip shape 4 in a base material 5, and is schematically shown in FIG. It was confirmed that high arrestability was obtained when the short crack branch 3a as shown was confirmed. It is presumed that the stress is relieved by the crack branch 3a.
2. In order to obtain the above fractured surface form, it is necessary to make the structure form to branch the crack. Here, the steel structure mainly composed of bainite in which a packet or the like is present is more advantageous than the steel structure mainly composed of ferrite, and the (100) surface which is a cleavage plane is propagated by cracks. It is effective to accumulate at an angle with respect to the rolling direction or the sheet width direction.
3. On the other hand, if the degree of integration on the (100) plane is too high, a large crack branch is generated from a very short crack branch. As described in Non-Patent Document 2 showing the design guideline for brittle crack arrest of ship structure, it is necessary to suppress the branching of brittle cracks in the standard ESSO test. It is necessary to specify the upper limit of the surface integration degree.
4). As a result of detailed observation and analysis of the fracture surface of the standard ESSO test, it is effective to improve the arrest performance by controlling the material at the center of the plate thickness, which is the tip of the crack. It is effective to satisfy the following equation (2) as an index related to the organization.
vTrs (1 / 2t) −12 × I RD // (110) [1 / 2t] ≦ −70 (2)
In the above formula (2),
vTrs (1 / 2t) : Fracture surface transition temperature (° C.) at the center of the plate thickness (= 1 / 2t)
I RD // (110) [1 / 2t] : The degree of integration of the RD // (110) plane of the central portion of the plate thickness (= 1 / 2t), where t is the plate thickness (mm).
5. Further, the structure is refined by rolling to a cumulative reduction rate of 20% or more in the state of the austenite recrystallization temperature range, and then the cumulative reduction rate in the state of the austenite non-recrystallization temperature range. The texture at the center of the plate thickness is controlled by performing rolling in which the difference between the rolling temperature of the first pass and the rolling temperature of the last pass is 40 ° C. or less by controlling the texture at the center of the plate thickness to 40 to 70%. Can be realized.
6). As a technique for improving the toughness of high heat input welds, the composite sulfides of TiN, CaS and MnS are finely divided to suppress grain growth when exposed to high temperatures of welding, and in the subsequent cooling process It is effective to promote internal transformation and refine the heat affected zone structure at room temperature.
　本発明は得られた知見を基に更に検討を加えてなされたものである。すなわち、本発明は、
１．鋼組成が、質量％で、Ｃ：０．０３～０．１５％、Ｓｉ：０．０１～０．５％、Ｍｎ：１．４０～２．５０％、Ａｌ：０．００５～０．０８％、Ｐ：０．０３％以下、Ｓ：０．０００５～０．００３０％、Ｎ：０．００３６～０．００７０％、Ｔｉ：０．００４～０．０３０％、Ｃａ：０．０００５～０．００３０％を含有し、且つ、Ｃａ、Ｓ、Ｏの各含有量が、下記（１）式を満足し、残部がＦｅおよび不可避的不純物で、金属組織がベイナイトを主体とし、板厚中央部におけるＲＤ／／（１１０）面の集積度が１．５～４．０の集合組織を有し、かつ表層部および板厚中央部におけるシャルピー破面遷移温度ｖＴｒｓが－４０℃以下であることを特徴とする脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。
０．３０≦（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓ≦０．８０・・・（１）
ただし、式（１）において、Ｃａ、Ｏ、Ｓは含有量（質量％）とする。
２．鋼組成が、更に、質量％で、Ｎｂ：０．０５％以下、Ｃｕ：１．０％以下、Ｎｉ：１．０％以下、Ｃｒ：０．５％以下、Ｍｏ：０．５％以下、Ｖ：０．２％以下、Ｂ：０．００３％以下、ＲＥＭ：０．０１％以下の１種または２種以上を含有することを特徴とする１に記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。
３．板厚中央部のシャルピー破面遷移温度およびＲＤ／／（１１０）面の集積度が、下記（２）式を満たすことを特徴とする１または２記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。
ただし、式（２）において、
ｖＴｒｓ（１／２ｔ）－１２×ＩＲＤ／／（１１０）［１／２ｔ］≦－７０・・・（２）
ｖＴｒｓ（１／２ｔ）：板厚中央部（１／２ｔ）の破面遷移温度（℃）
ＩＲＤ／／（１１０）［１／２ｔ］：板厚中央部（１／２ｔ）のＲＤ／／（１１０）面の集積度
とする。なお、ｔは板厚（ｍｍ）である。
４．１または２に記載の組成を有する鋼素材を、１０００～１２００℃の温度に加熱し、オーステナイト再結晶温度域およびオーステナイト未再結晶温度域における累積圧下率の合計が６５％以上の圧延を実施し、このとき、板厚中央部がオーステナイト再結晶温度域にある状態においては累積圧下率が２０％以上の圧延を行い、次いで、板厚中央部がオーステナイト未再結晶温度域にある状態においては、累積圧下率が４０～７０％とする圧延を行い、かつ、前記板厚中央部がオーステナイト未再結晶温度域にある状態における圧延のうち最初のパスの圧延温度と最後のパスの圧延温度との差が４０℃以内であり、その後、４．０℃／ｓ以上の冷却速度にて４５０℃以下まで冷却することを特徴とする脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板の製造方法。
５．４５０℃以下に加速冷却した後、さらに、Ａｃ１点以下の温度に焼戻す工程を有することを特徴とする４に記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板の製造方法。 The present invention has been made by further study based on the obtained knowledge. That is, the present invention
1. Steel composition is mass%, C: 0.03-0.15%, Si: 0.01-0.5%, Mn: 1.40-2.50%, Al: 0.005-0.08 %, P: 0.03% or less, S: 0.0005 to 0.0030%, N: 0.0036 to 0.0070%, Ti: 0.004 to 0.030%, Ca: 0.0005 to 0 .0030%, and each content of Ca, S, and O satisfies the following formula (1), the balance is Fe and inevitable impurities, the metal structure is mainly bainite, and the central portion of the plate thickness The RD // (110) plane has a texture of 1.5 to 4.0, and the Charpy fracture surface transition temperature vTrs in the surface layer portion and the thickness center portion is −40 ° C. or lower. High strength thick steel plate for high heat input welding with excellent brittle crack propagation stopping characteristics.
0.30 ≦ (Ca− (0.18 + 130 × Ca) × O) /1.25/S≦0.80 (1)
However, in Formula (1), Ca, O, and S are made into content (mass%).
2. The steel composition is further mass%, Nb: 0.05% or less, Cu: 1.0% or less, Ni: 1.0% or less, Cr: 0.5% or less, Mo: 0.5% or less, V: 0.2% or less, B: 0.003% or less, REM: 0.01% or less of one type or two or more types, which is excellent in brittle crack propagation stop property according to 1. High strength thick steel plate for high heat input welding.
3. The large insertion with excellent brittle crack propagation stop characteristics according to 1 or 2, wherein the Charpy fracture transition temperature at the center of the plate thickness and the degree of integration of the RD // (110) plane satisfy the following formula (2): High strength thick steel plate for thermal welding.
However, in Formula (2),
vTrs (1 / 2t) −12 × I RD // (110) [1 / 2t] ≦ −70 (2)
vTrs (1 / 2t) : Fracture surface transition temperature (° C.) at the center of the plate thickness (1 / 2t)
I RD // (110) [1 / 2t] : The degree of integration of the RD // (110) plane in the central portion (1 / 2t) of the plate thickness. Here, t is a plate thickness (mm).
The steel material having the composition described in 4.1 or 2 is heated to a temperature of 1000 to 1200 ° C., and rolling with a total cumulative rolling reduction of 65% or more in the austenite recrystallization temperature range and the austenite non-recrystallization temperature range is performed. At this time, in the state where the plate thickness central portion is in the austenite recrystallization temperature range, the rolling reduction is performed at a cumulative reduction ratio of 20% or more, and then in the state where the plate thickness central portion is in the austenite non-recrystallization temperature range. Is the rolling temperature of the first pass and the rolling temperature of the last pass of the rolling in a state where the rolling reduction is 40 to 70% and the central part of the plate thickness is in the austenite non-recrystallization temperature range. The heat input is excellent in brittle crack propagation stopping characteristics, characterized in that the difference is within 40 ° C. and then cooled to 450 ° C. or lower at a cooling rate of 4.0 ° C./s or higher. Method of producing a high strength steel plate for contact.
5. After accelerated cooling to 450 ° C. or lower, and further tempering to a temperature of A c1 point or lower, the high strength thickness for high heat input welding having excellent brittle crack propagation stopping characteristics according to 4 A method of manufacturing a steel sheet.
　本発明によれば、板厚方向に集合組織が適切に制御され、脆性亀裂伝播停止特性、大入熱溶接継手靭性に優れる、高強度厚肉鋼板およびその製造方法が得られる。板厚５０ｍｍ以上、好ましくは板厚５０ｍｍ超え、より好ましくは板厚５５ｍｍ以上、一層好ましくは板厚６０ｍｍ以上の鋼板に本発明を適用することが、従来技術に係る鋼に対してより顕著な優位性を発揮するため、有効である。そして、例えば、造船分野では大型のコンテナ船、バルクキャリアーの強力甲板部構造においてハッチサイドコーミングや甲板部材へ本発明を適用することにより船舶の安全性向上に寄与するなど、産業上極めて有用である。 According to the present invention, it is possible to obtain a high-strength thick steel plate and a method for producing the same, in which the texture is appropriately controlled in the plate thickness direction, and the brittle crack propagation stop property and the high heat input weld joint toughness are excellent. Applying the present invention to a steel plate having a plate thickness of 50 mm or more, preferably more than 50 mm, more preferably 55 mm or more, and even more preferably 60 mm or more is more significant than the steel according to the prior art. It is effective because it demonstrates its properties. And, for example, in the shipbuilding field, it contributes to improving the safety of ships by applying the present invention to hatch side combing and deck members in the structure of large container ships and bulk carrier strong deck parts. .
図１は、板厚５０ｍｍを超える厚鋼板の標準ＥＳＳＯ試験の破面形態を模式的に示す図であり、（ａ）は試験片を平面側から観察した図、（ｂ）は試験片の破面を示す図である。FIG. 1 is a diagram schematically showing a fracture surface form of a standard ESSO test of a thick steel plate having a thickness of more than 50 mm, (a) is a view of the test piece observed from the plane side, and (b) is a fracture of the test piece. It is a figure which shows a surface.
　本発明では、１．鋼組成、２．板厚表層部および中央部の靭性と板厚中央部の集合組織、３．金属組織、および４．製造条件を規定する。 In the present invention, 1. Steel composition, 2. 2. the toughness of the surface thickness layer and the central portion and the texture of the central portion of the thickness; 3. metal structure, and Define manufacturing conditions.
　１．鋼組成
　以下、本発明における好ましい化学成分について説明する。説明において％は質量％とする。 1. Steel composition Hereinafter, preferred chemical components in the present invention will be described. In the description,% is mass%.
　Ｃ：０．０３～０．１５％
Ｃは鋼の強度を向上する元素であり、本発明では、所望の強度を確保するためには０．０３％以上の含有を必要とする。一方、０．１５％を超えると、溶接性が劣化するばかりか靭性にも悪影響がある。このため、Ｃは、０．０３～０．１５％の範囲に規定する。好ましくは、０．０５～０．１５％である。 C: 0.03-0.15%
C is an element that improves the strength of steel, and in the present invention, it is necessary to contain 0.03% or more in order to ensure a desired strength. On the other hand, if it exceeds 0.15%, the weldability is deteriorated and the toughness is also adversely affected. Therefore, C is specified in the range of 0.03 to 0.15%. Preferably, it is 0.05 to 0.15%.
　Ｓｉ：０．０１～０．５％
Ｓｉは脱酸元素として、また、鋼の強化元素として有効である。しかし、０．０１％未満の含有量ではその効果がない。一方、０．５％を超えると鋼の表面性状を損なうばかりか靭性が極端に劣化する。従ってその添加量を０．０１～０．５％とする。好ましくは、０．０２～０．４５％の範囲である。 Si: 0.01 to 0.5%
Si is effective as a deoxidizing element and as a steel strengthening element. However, when the content is less than 0.01%, the effect is not obtained. On the other hand, if it exceeds 0.5%, not only the surface properties of the steel are impaired, but also the toughness is extremely deteriorated. Therefore, the addition amount is set to 0.01 to 0.5%. Preferably, it is 0.02 to 0.45% of range.
　Ｍｎ：１．４０～２．５０％
Ｍｎは、強化元素として添加する。１．４０％より少ないとその効果が十分でない。一方、２．５０％を超えると溶接性が劣化し、鋼材コストも上昇する。このため、Ｍｎは１．４０～２．５０％とする。好ましくは、１．４２～２．４０％の範囲である。 Mn: 1.40-2.50%
Mn is added as a strengthening element. If less than 1.40%, the effect is not sufficient. On the other hand, if it exceeds 2.50%, the weldability deteriorates and the steel material cost also increases. Therefore, Mn is set to 1.40 to 2.50%. Preferably, it is in the range of 1.42 to 2.40%.
　Ｐ：０．０３％以下
Ｐは、０．０３％を超えると、溶接部の靭性を著しく劣化させる。このため上限を０．０３％とする。好ましくは０．０２％以下である。 P: 0.03% or less When P exceeds 0.03%, the toughness of the welded portion is significantly deteriorated. For this reason, the upper limit is made 0.03%. Preferably it is 0.02% or less.
　Ｓ：０．０００５～０．００３０％
Ｓは、必要なＣａＳおよびＭｎＳを生成させるために０．０００５％以上必要である。一方、０．００３０％を超えると母材の靭性を劣化させる。このため、Ｓは０．０００５～０．００３０％とする。好ましくは、０．０００６～０．００２５％の範囲である。 S: 0.0005 to 0.0030%
S is required to be 0.0005% or more in order to generate necessary CaS and MnS. On the other hand, if it exceeds 0.0030%, the toughness of the base material is deteriorated. Therefore, S is set to 0.0005 to 0.0030%. Preferably, it is in the range of 0.0006 to 0.0025%.
　Ａｌ：０．００５～０．０８％
Ａｌは、脱酸剤として作用し、このためには０．００５％以上の含有を必要とする。しかし、０．０８％を超えて含有すると、靭性を低下させるとともに、溶接した場合に、溶接金属部の靭性を低下させる。このため、Ａｌは、０．００５～０．０８％の範囲に規定する。好ましくは、０．０２～０．０６％である。 Al: 0.005 to 0.08%
Al acts as a deoxidizing agent, and for this purpose, a content of 0.005% or more is required. However, when it contains exceeding 0.08%, while reducing toughness, when welding, the toughness of a weld metal part will be reduced. For this reason, Al is specified in the range of 0.005 to 0.08%. Preferably, it is 0.02 to 0.06%.
　Ｔｉ：０．００４～０．０３０％
Ｔｉは微量の添加により、窒化物、炭化物、あるいは炭窒化物を形成し、溶接熱影響部でのオーステナイトの粗大化を抑制することにより、および／または、フェライト変態核としてフェライト変態を促進することにより、結晶粒を微細化して母材靭性を向上させる効果を有する。その効果は０．００４％以上の添加によって得られる。しかし、０．０３０％を超える含有は、ＴｉＮ粒子の粗大化により、母材および溶接熱影響部の靭性を低下させる。このため、Ｔｉは、０．００４～０．０３０％の範囲にする。好ましくは、０．００６～０．０２８％の範囲である。 Ti: 0.004 to 0.030%
Ti can form nitrides, carbides, or carbonitrides by adding a small amount, suppress austenite coarsening in the weld heat affected zone, and / or promote ferrite transformation as a ferrite transformation nucleus. This has the effect of refining the crystal grains and improving the base material toughness. The effect is obtained by adding 0.004% or more. However, the content exceeding 0.030% reduces the toughness of the base material and the weld heat affected zone due to the coarsening of the TiN particles. Therefore, Ti is set in the range of 0.004 to 0.030%. Preferably, it is 0.006 to 0.028% of range.
　Ｎ：０．００３６～０．００７０％
Ｎは、ＴｉＮの必要量を確保する上で必要な元素である。０．００３６％未満では十分なＴｉＮ量が得られず、溶接部靭性が劣化する。０．００７０％を超えると、溶接熱サイクルを受けた際にＴｉＮが再固溶して固溶Ｎが過剰に生成して靭性が著しく劣化する。このため、Ｎは０．００３６～０．００７０％とする。好ましくは、０．００３８～０．００６５％の範囲である。 N: 0.0036 to 0.0070%
N is an element necessary for securing the necessary amount of TiN. If it is less than 0.0036%, a sufficient amount of TiN cannot be obtained, and the weld toughness deteriorates. If it exceeds 0.0070%, TiN will re-dissolve when subjected to the welding heat cycle, and excessive N will be generated and the toughness will deteriorate significantly. For this reason, N is made 0.0036 to 0.0070%. Preferably, it is 0.0038 to 0.0065% of range.
　Ｃａ：０．０００５～０．００３０％
Ｃａは、Ｓの固定による靭性改善効果を有する元素である。このような効果を発揮させるには少なくとも０．０００５％以上含有する必要がある。しかし、０．００３０％を超えて含有しても効果が飽和する。このため、本発明では、Ｃａは０．０００５～０．００３０％の範囲に限定する。好ましくは、０．０００７～０．００２８％の範囲である。 Ca: 0.0005 to 0.0030%
Ca is an element having an effect of improving toughness by fixing S. In order to exhibit such an effect, it is necessary to contain at least 0.0005% or more. However, the effect is saturated even if the content exceeds 0.0030%. Therefore, in the present invention, Ca is limited to the range of 0.0005 to 0.0030%. Preferably, it is in the range of 0.0007 to 0.0028%.
　本発明において、以下の式（１）を満足する必要がある。
０．３０≦（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓ≦０．８０・・・（１）
ただし、式（１）において、Ｃａ、Ｏ、Ｓは含有量（質量％）とする。
ＣａおよびＳは、（１）式の関係を満足するように含有させる必要がある。この場合には、ＣａＳ上にＭｎＳが析出した複合硫化物の形態となる。この複合硫化物がフェライト変態の核として機能するので、溶接熱影響部の組織が微細化され、溶接熱影響部の靭性が向上する。（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓの値が０．３０に満たないと、ＣａＳが晶出しないためにＳはＭｎＳ単独の形態で析出する。このＭｎＳは鋼板製造時の圧延で伸長されて母材靭性低下を引き起こすとともに、本発明の主眼である溶接熱影響部でＭｎＳが溶融するために微細分散が達成されない。一方、（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓの値が０．８０を超えると、ＳがほとんどＣａによって固定され、フェライト生成核として作用するＭｎＳがＣａＳ上に析出しないため、十分な靭性向上が達成されない。（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓの値の好ましい範囲は、０．３２～０．７８％である。 In the present invention, it is necessary to satisfy the following formula (1).
0.30 ≦ (Ca− (0.18 + 130 × Ca) × O) /1.25/S≦0.80 (1)
However, in Formula (1), Ca, O, and S are made into content (mass%).
Ca and S must be contained so as to satisfy the relationship of the formula (1). In this case, it becomes the form of the composite sulfide in which MnS is deposited on CaS. Since this composite sulfide functions as a nucleus of ferrite transformation, the structure of the weld heat affected zone is refined, and the toughness of the weld heat affected zone is improved. If the value of (Ca− (0.18 + 130 × Ca) × O) /1.25/S is less than 0.30, since CaS does not crystallize, S precipitates in the form of MnS alone. This MnS is elongated by rolling at the time of manufacturing the steel sheet to cause a reduction in the toughness of the base material, and fine dispersion is not achieved because MnS melts in the weld heat affected zone which is the main point of the present invention. On the other hand, when the value of (Ca− (0.18 + 130 × Ca) × O) /1.25/S exceeds 0.80, S is almost fixed by Ca, and MnS acting as ferrite nuclei is formed on CaS. Since it does not precipitate, sufficient toughness improvement is not achieved. A preferable range of the value of (Ca− (0.18 + 130 × Ca) × O) /1.25/S is 0.32 to 0.78%.
　以上が本発明の基本成分組成である。更に特性を向上させるため、Ｎｂ、Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏ、Ｖ、Ｂ、ＲＥＭの１種以上を含有することが可能である。 The above is the basic component composition of the present invention. In order to further improve the characteristics, it is possible to contain one or more of Nb, Cu, Ni, Cr, Mo, V, B, and REM.
　Ｎｂ：０．０５％以下
Ｎｂは、ＮｂＣとしてフェライト変態時あるいは再加熱時に析出し、高強度化に寄与する。また、オーステナイト域の圧延において未再結晶温度域を拡大させる効果をもちベイナイトのパケットの細粒化に寄与するので、靭性の改善にも有効である。その効果は０．００５％以上含有することにより発揮されるので、含有させる場合には、０．００５％以上とすることが好ましい。しかしながら、０．０５％を超えて添加すると、粗大なＮｂＣが析出し、逆に靭性の低下を招くので、含有させる場合には、その上限を０．０５％とするのが好ましい。より好ましくは、０．００７～０．０４５％の範囲である。 Nb: 0.05% or less Nb precipitates as NbC during ferrite transformation or reheating, and contributes to increasing the strength. In addition, it has the effect of expanding the non-recrystallization temperature range in rolling in the austenite range, and contributes to the fine graining of bainite packets, so it is also effective in improving toughness. Since the effect is exhibited by containing 0.005% or more, when it contains, it is preferable to make it 0.005% or more. However, if added over 0.05%, coarse NbC precipitates and conversely causes a decrease in toughness. When it is contained, the upper limit is preferably made 0.05%. More preferably, it is in the range of 0.007 to 0.045%.
　Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏ
Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏはいずれも鋼の焼入れ性を高める元素である。圧延後の強度アップに直接寄与するとともに、靭性、高温強度、あるいは耐候性などの機能向上のために添加することができ、これらの効果は０．０１％以上含有することにより発揮されるので、含有させる場合には、０．０１％以上とすることが好ましい。しかしながら、過度に含有すると靭性や溶接性が劣化するため、含有させる場合には、それぞれ上限をＣｕは１．０％、Ｎｉは１．０％、Ｃｒは０．５％、Ｍｏは０．５％とすることが好ましい。より好ましくは、Ｃｕ：０．０２～０．９５％、Ｎｉ：０．０２～０．９５％、Ｃｒ：０．０２～０．４６％、Ｍｏ：０．０２～０．４６％の範囲である。 Cu, Ni, Cr, Mo
Cu, Ni, Cr, and Mo are all elements that enhance the hardenability of steel. While contributing directly to strength enhancement after rolling, it can be added to improve functions such as toughness, high-temperature strength, or weather resistance, since these effects are exhibited by containing 0.01% or more, When contained, the content is preferably 0.01% or more. However, if it is excessively contained, toughness and weldability are deteriorated. Therefore, when it is included, the upper limit is 1.0% for Cu, 1.0% for Ni, 0.5% for Cr, and 0.5% for Mo. % Is preferable. More preferably, Cu: 0.02 to 0.95%, Ni: 0.02 to 0.95%, Cr: 0.02 to 0.46%, Mo: 0.02 to 0.46% is there.
　Ｖ：０．２％以下
Ｖは、Ｖ（Ｃ、Ｎ）として析出強化により、鋼の強度を向上する元素であり、この効果を発揮させるために０．００１％以上含有させてもよい。しかし、０．２％を超えて含有すると、靭性を低下させる。このため、Ｖを含有させる場合には、０．２％以下とすることが好ましく、０．００１～０．１０％の範囲とすることがより好ましい。 V: 0.2% or less V is an element that improves the strength of steel by precipitation strengthening as V (C, N), and may be contained by 0.001% or more in order to exert this effect. However, when it contains exceeding 0.2%, toughness will be reduced. Therefore, when V is contained, the content is preferably 0.2% or less, and more preferably in the range of 0.001 to 0.10%.
　Ｂ：０．００３％以下
Ｂは微量で鋼の焼き入れ性を高める元素であり、この効果を発揮させるために０．０００５％以上含有させてもよい。しかし、０．００３％を超えて含有すると溶接部の靭性を低下させるので、Ｂを含有させる場合には０．００３％以下とすることが好ましい。より好ましくは、０．０００６～０．００２５％の範囲である。 B: 0.003% or less B is an element that enhances the hardenability of steel in a small amount, and may be contained by 0.0005% or more in order to exert this effect. However, if it exceeds 0.003%, the toughness of the welded portion is lowered. Therefore, when B is contained, the content is preferably 0.003% or less. More preferably, it is in the range of 0.0006 to 0.0025%.
　ＲＥＭ：０．０１％以下
ＲＥＭは溶接熱影響部の組織を微細化し靭性を向上させ、添加しても本発明の効果が損なわれることはないので必要に応じて添加してもよい。この効果は０．００１０％以上含有することにより発揮されるので、含有させる場合には、０．００１０％以上とすることが好ましい。しかし、過度に添加すると、粗大な介在物を形成し母材の靭性を劣化させるので、添加する場合には、添加量の上限を０．０１％とするのが好ましい。 REM: 0.01% or less REM refines the structure of the weld heat-affected zone to improve toughness, and even if added, the effect of the present invention is not impaired, so it may be added as necessary. Since this effect is exhibited by containing 0.0010% or more, when it is contained, the content is preferably 0.0010% or more. However, if added excessively, coarse inclusions are formed and the toughness of the base material is deteriorated. Therefore, when added, the upper limit of the addition amount is preferably 0.01%.
　なお、Ｏは不可避的不純物として鋼中に含有され、清浄度を低下させる。このため、本発明ではできるだけＯを低減することが望ましい。特に、Ｏ含有量が０．００５０％を超えるとＣａＯ系介在物が粗大化して母材靭性を低下させてしまう。このため、好ましくは０．００５０％以下とする。 In addition, O is contained in steel as an unavoidable impurity and reduces cleanliness. For this reason, in the present invention, it is desirable to reduce O as much as possible. In particular, when the O content exceeds 0.0050%, CaO inclusions are coarsened and the base material toughness is lowered. For this reason, Preferably it is 0.0050% or less.
　本発明では、ＣａをＣａＳとして晶出させるために、Ｃａと結合力の強いＯ量をＣａ添加前に低減させておくことが必要であり、Ｃａ添加前の残存酸素量は、０．００５０％以下であることが好ましい。残存酸素量の低減方法としては、脱ガスを強化する、あるいは、脱酸剤を投入する、などの方法をとることができる。 In the present invention, in order to crystallize Ca as CaS, it is necessary to reduce the amount of O having strong binding force with Ca before adding Ca, and the residual oxygen amount before adding Ca is 0.0050%. The following is preferable. As a method for reducing the amount of residual oxygen, a method such as enhancing degassing or introducing a deoxidizer can be employed.
　上記した成分以外の残部は、Ｆｅおよび不可避的不純物である。 The balance other than the above components is Fe and inevitable impurities.
　２．板厚表層部および中央部の靭性と板厚中央部の集合組織
　本発明では、圧延方向または圧延直角方向など水平方向（鋼板の面内方向）に進展する亀裂に対して脆性亀裂伝播停止特性を向上させるため、板厚表層部および中央部での靭性と、板厚中央部におけるＲＤ／／（１００）面の集積度とを、所望する脆性亀裂伝播停止特性に応じて適宜規定する。 2. In the present invention, the brittle crack propagation stop property is exhibited for cracks that develop in the horizontal direction (in-plane direction of the steel sheet) such as the rolling direction or the direction perpendicular to the rolling direction. In order to improve, the toughness at the surface thickness layer and the central portion and the degree of integration of the RD // (100) plane at the central portion of the thickness are appropriately defined according to the desired brittle crack propagation stop characteristics.
　まず、母材靭性が良好であることが亀裂の進展を抑制するための前提となる。本発明に係る鋼板では、板厚表層部および中央部での靭性として、板厚表層部および中央部におけるシャルピー破面遷移温度を－４０℃以下と規定する。なお、板厚中央部におけるシャルピー破面遷移温度は－５０℃以下であることが好ましい。 First, good base material toughness is a precondition for suppressing the progress of cracks. In the steel sheet according to the present invention, the Charpy fracture surface transition temperature at the plate thickness surface layer portion and the central portion is defined as −40 ° C. or less as the toughness at the plate thickness surface layer portion and the center portion. The Charpy fracture surface transition temperature at the center of the plate thickness is preferably −50 ° C. or lower.
　また、ＲＤ／／（１００）面の集合組織を発達させることにより、へき開面を亀裂主方向に対し斜めに集積させ、微細な亀裂分岐を発生させることによる脆性亀裂先端の応力緩和の効果により脆性亀裂伝播停止性能が向上する。最近のコンテナ船やバルクキャリアーなど船体外板に用いられるようになった板厚５０ｍｍを超える厚肉材で、構造安全性を確保する上で目標とされる脆性亀裂伝播停止性能：Ｋｃａ（－１０℃）≧６０００Ｎ／ｍｍ３／２を得る場合、板厚中央部におけるＲＤ／／（１１０）面の集積度を１．５以上、好ましくは１．７以上とする必要がある。したがって、本発明では、板厚中央部におけるＲＤ／／（１１０）面の集積度を１．５以上、好ましくは１．７以上とする。 In addition, by developing a texture of the RD // (100) plane, the cleavage plane is accumulated obliquely with respect to the main crack direction, and the effect of stress relaxation at the brittle crack tip by generating fine crack branching causes brittleness. The crack propagation stop performance is improved. A brittle crack propagation stopping performance targeted for ensuring structural safety: Kca (−10), which is a thick material exceeding 50 mm thick that has been used for hull outer plates such as recent container ships and bulk carriers. In order to obtain (° C.) ≧ 6000 N / mm 3/2 , the degree of integration of the RD // (110) plane in the central portion of the plate thickness needs to be 1.5 or more, preferably 1.7 or more. Therefore, in the present invention, the degree of integration of the RD // (110) plane at the center of the plate thickness is 1.5 or more, preferably 1.7 or more.
　一方、板厚中央部におけるＲＤ／／（１１０）面の集積度が４．０を超えると集合組織が過度に発達するため、微細な亀裂分岐が発生するのではなく脆性亀裂が明瞭に分岐してしまうため、脆性亀裂先端の応力緩和の効果による脆性亀裂伝播停止性能が発揮されにくくなる。このため、ＲＤ／／（１１０）面の集積度を１．５～４．０の範囲とする。 On the other hand, when the degree of integration of the RD // (110) plane at the center of the plate thickness exceeds 4.0, the texture develops excessively, so that a brittle crack is clearly branched rather than a fine crack branching. Therefore, the brittle crack propagation stopping performance due to the stress relaxation effect at the tip of the brittle crack becomes difficult to be exhibited. For this reason, the integration degree of the RD // (110) plane is set to a range of 1.5 to 4.0.
　ここで、板厚中央部におけるＲＤ／／（１１０）面の集積度とは、次のことを指す。まず、板厚中央部から板厚１ｍｍのサンプルを採取し、板面に平行な面を機械研磨・電解研磨することにより、Ｘ線回折用の試験片を用意する。この試験片を用いて、Ｍｏ線源を用いてＸ線回折装置を使用して、Ｘ線回折測定を実施し、（２００）、（１１０）および（２１１）正極点図を求め、得られた正極点図から３次元結晶方位密度関数をＢｕｎｇｅ法で計算して求める。次に、得られた３次元結晶方位密度関数から、Ｂｕｎｇｅ表記でψ２＝０°～９０°まで、５°間隔で合計１９枚の断面図において、圧延方向に対して（１１０）面が平行となる方位の３次元結晶方位密度関数の値を積算して積算値を求め、この積算値を前記積算した方位の個数で割った値を、ＲＤ／／（１１０）面の集積度と称する。 Here, the degree of integration of the RD // (110) plane in the central portion of the plate thickness refers to the following. First, a sample with a plate thickness of 1 mm is taken from the center of the plate thickness, and a test piece for X-ray diffraction is prepared by mechanically polishing and electrolytic polishing a surface parallel to the plate surface. Using this test piece, an X-ray diffraction measurement was performed using an X-ray diffractometer using a Mo ray source, and (200), (110) and (211) positive electrode dot diagrams were obtained and obtained. A three-dimensional crystal orientation density function is calculated from the positive electrode dot diagram by the Bunge method. Next, from the obtained three-dimensional crystal orientation density function, the (110) plane is parallel to the rolling direction in a total of 19 cross-sectional views at 5 ° intervals from ψ 2 = 0 ° to 90 ° in Bunge notation. The integrated value is obtained by integrating the values of the three-dimensional crystal orientation density function of the orientation, and the value obtained by dividing the integrated value by the number of the integrated orientations is referred to as the degree of integration of the RD // (110) plane.
　上述の母材靭性および集合組織の規定に加えて、板厚中央部のシャルピー破面遷移温度およびＲＤ／／（１１０）面の集積度が、下記（２）式を満たすことが、好ましい。下記（２）式が満足されることにより、さらに優れた脆性亀裂伝播停止性能を得ることができる。
ｖＴｒｓ（１／２ｔ）－１２×ＩＲＤ／／（１１０）［１／２ｔ］≦－７０・・・（２）
ただし、式（２）において
ｖＴｒｓ（１／２ｔ）：板厚中央部のシャルピー破面遷移温度（℃）
ＩＲＤ／／（１１０）［１／２ｔ］：板厚中央部のＲＤ／／（１１０）集積度
とする。なお、ｔは板厚（ｍｍ）である。 It is preferable that the Charpy fracture surface transition temperature at the center of the plate thickness and the degree of integration of the RD // (110) plane satisfy the following formula (2) in addition to the above-mentioned provisions of the base material toughness and texture. When the following expression (2) is satisfied, further excellent brittle crack propagation stopping performance can be obtained.
vTrs (1 / 2t) −12 × I RD // (110) [1 / 2t] ≦ −70 (2)
However, in formula (2), vTrs (1 / 2t) : Charpy fracture surface transition temperature (° C.) at the center of the plate thickness
I RD // (110) [1 / 2t] : RD // (110) integration degree at the center of the plate thickness. Here, t is a plate thickness (mm).
　３．金属組織
　上記の靭性および集合組織を得るためには、オーステナイト未再結晶温度域において制御圧延を行った後に、ベイナイトへ変態させることが有効である。圧延後にオーステナイトからフェライトへ変態する場合は、目的とする靭性は得られるものの、オーステナイトからフェライトへ変態する際に変態時間が十分に存在するため、得られる集合組織がランダムとなってしまい、目標とするＲＤ／／（１１０）面の集積度が１．５以上、好ましくは１．７以上、が達成できない。これに対して、オーステナイト未再結晶温度域で圧延された組織がベイナイトへ変態する場合は変態時間が十分ではなく、特定方位の集合組織が優先的に形成される、いわゆるバリアントの選択が行われることにより、ＲＤ／／（１１０）面の集積度が１．５以上、好ましくは１．７以上、を得ることができる。このため圧延・冷却後に得られる金属組織はベイナイト主体とする。本発明で、金属組織がベイナイト主体であるとは、ベイナイト相の面積分率が全体の８０％以上であることとする。残部は、フェライト、マルテンサイト（島状マルテンサイトを含む）、パーライトなどが合計の面積分率で２０％以下であれば許容される。 3. Metal structure In order to obtain the above toughness and texture, it is effective to transform into bainite after performing controlled rolling in the austenite non-recrystallization temperature range. When transforming from austenite to ferrite after rolling, the desired toughness can be obtained, but because there is sufficient transformation time when transforming from austenite to ferrite, the resulting texture becomes random, and the target The degree of integration of the RD // (110) plane cannot be 1.5 or more, preferably 1.7 or more. On the other hand, when the structure rolled in the austenite non-recrystallization temperature region transforms into bainite, the transformation time is not sufficient, and a so-called variant is selected in which a texture with a specific orientation is preferentially formed. As a result, the degree of integration of the RD // (110) plane can be 1.5 or more, preferably 1.7 or more. For this reason, the metal structure obtained after rolling and cooling is mainly bainite. In the present invention, that the metal structure is mainly bainite is that the area fraction of the bainite phase is 80% or more of the whole. The balance is acceptable if ferrite, martensite (including island martensite), pearlite, etc. are 20% or less in total area fraction.
　４．製造条件
　以下、本発明における好ましい製造条件について説明する。 4). Manufacturing conditions Hereinafter, preferable manufacturing conditions in the present invention will be described.
　製造条件としては、鋼素材の加熱温度、熱間圧延条件、冷却条件などを規定することが好ましい。特に、熱間圧延については、オーステナイト再結晶温度域およびオーステナイト未再結晶温度域の合計での累積圧下率のほかに、板厚中央部がオーステナイト再結晶温度域にある場合と、オーステナイト未再結晶温度域にある場合とのそれぞれについて、累積圧下率を規定するとともに、板厚中央部がオーステナイト未再結晶域にある状態における圧延の温度条件を規定することが好ましい。これらを規定することにより、厚鋼板の表層部および板厚中央部におけるシャルピー破面遷移温度ｖＴｒｓ、板厚中央部におけるＲＤ／／（１１０）集積度を、所望の値とすることができる。 As manufacturing conditions, it is preferable to prescribe the heating temperature, hot rolling conditions, cooling conditions, etc. of the steel material. In particular, for hot rolling, in addition to the cumulative reduction ratio in the sum of the austenite recrystallization temperature range and the austenite non-recrystallization temperature range, the case where the central portion of the plate thickness is in the austenite recrystallization temperature range, It is preferable to define the cumulative rolling reduction for each of the cases in the temperature range and the rolling temperature conditions in a state where the central portion of the plate thickness is in the austenite non-recrystallized region. By defining these, the Charpy fracture surface transition temperature vTrs in the surface layer portion and the plate thickness center portion of the thick steel plate, and the RD // (110) integration degree in the plate thickness center portion can be set to desired values.
　まず、上記した組成の溶鋼を、転炉等で溶製し、連続鋳造等で鋼素材（スラブ）とする。 First, molten steel having the above composition is melted in a converter or the like, and is made into a steel material (slab) by continuous casting or the like.
　ついで、鋼素材を、１０００～１２００℃の温度に加熱してから熱間圧延を行うことが好ましい。加熱温度が１０００℃未満では、オーステナイト再結晶温度域における圧延を行う時間が十分に確保できない。また、１２００℃超えではオーステナイト粒が粗大化し、靭性の低下を招くばかりか、酸化ロスが顕著となり、歩留が低下する。したがって、加熱温度は１０００～１２００℃とすることが好ましい。靭性の観点からより好ましい加熱温度の範囲は１０００～１１５０℃である。 Next, it is preferable to perform hot rolling after heating the steel material to a temperature of 1000 to 1200 ° C. If the heating temperature is less than 1000 ° C., sufficient time for rolling in the austenite recrystallization temperature region cannot be secured. When the temperature exceeds 1200 ° C., austenite grains become coarse, resulting in a decrease in toughness, and an oxidation loss becomes remarkable, resulting in a decrease in yield. Therefore, the heating temperature is preferably 1000 to 1200 ° C. A more preferable heating temperature range is 1000 to 1150 ° C. from the viewpoint of toughness.
　本発明においては、以下に述べるように、熱間圧延条件およびそれに続く冷却条件を規定することが好ましい。これにより、オーステナイト未再結晶温度域で圧延された組織をベイナイトへ変態させるので、この場合の変態時間が十分ではないことから、特定方位の集合組織が優先的に形成される、いわゆるバリアントの選択(variant selection)が行われることにより、ＲＤ／／（１１０）面の集積度を１．５以上、好ましくは１．７以上とすることができる。 In the present invention, as described below, it is preferable to define hot rolling conditions and subsequent cooling conditions. As a result, the structure rolled in the austenite non-recrystallization temperature region is transformed into bainite, so the transformation time in this case is not sufficient, so the so-called variant selection in which a texture with a specific orientation is preferentially formed. By performing (variant selection), the degree of integration of the RD // (110) plane can be 1.5 or more, preferably 1.7 or more.
　熱間圧延は、まず、板厚中央部がオーステナイト再結晶温度域にある状態において、累積圧下率を２０％以上とする圧延を行うことが好ましい。この累積圧下率を２０％以上とすることによりオーステナイトが細粒化し、最終的に得られる金属組織も細粒化して、靭性が向上する。累積圧下率が２０％未満であると、オーステナイトの細粒化が不十分で、最終的に得られる組織において靭性が向上しない。 In the hot rolling, first, it is preferable to perform rolling with a cumulative reduction ratio of 20% or more in a state where the central portion of the plate thickness is in the austenite recrystallization temperature region. By setting the cumulative rolling reduction to 20% or more, austenite is refined and the finally obtained metal structure is also refined to improve toughness. If the cumulative rolling reduction is less than 20%, austenite is not sufficiently refined and the toughness is not improved in the finally obtained structure.
　次に、板厚中央部の温度がオーステナイト未再結晶温度域にある状態において累積圧下率４０～７０％以上とする圧延を行うことが好ましい。この温度域での累積圧下率を４０％以上とすることにより、板厚中央部の集合組織を十分に発達させ、板厚中央部のＲＤ／／（１１０）面の集積度を１．５以上、好ましくは１．７以上とすることができる。 Next, it is preferable to perform rolling to a cumulative reduction ratio of 40 to 70% or more in a state where the temperature at the center of the plate thickness is in the austenite non-recrystallization temperature range. By setting the cumulative reduction ratio in this temperature range to 40% or more, the texture at the center of the plate thickness is sufficiently developed, and the degree of integration of the RD // (110) plane at the center of the plate thickness is 1.5 or more. , And preferably 1.7 or more.
　また、この温度域での累積圧下率が７０％を超えると、集合組織が過度に発達し、ＲＤ／／（１１０）面の集積度が４．０を超える。このため、累積圧下率の範囲を４０～７０％とする。 In addition, when the cumulative rolling reduction in this temperature range exceeds 70%, the texture is excessively developed, and the degree of accumulation on the RD // (110) plane exceeds 4.0. For this reason, the range of the cumulative rolling reduction is set to 40 to 70%.
　なお、板厚中央部の温度がオーステナイト未再結晶温度域にある状態における圧延に時間がかかり過ぎると組織が粗大化してしまい、靭性の低下をまねいてしまう。そのため、前記板厚中央部がオーステナイト未再結晶域にある状態における圧延のうち最初のパスの圧延温度と最後のパスの圧延温度との差を４０℃以内とすることが好ましい。ここで、圧延温度とは、圧延直前の鋼材の板厚中央部の温度を指す。板厚中央部の温度は、板厚、表面温度および熱履歴等から、シミュレーション計算等により求められる。例えば、差分法を用い、板厚方向の温度分布を計算することにより、鋼板の板厚中央部の温度が求められる。 In addition, if it takes too much time for rolling in a state where the temperature at the center of the plate thickness is in the austenite non-recrystallization temperature range, the structure becomes coarse and the toughness is lowered. Therefore, it is preferable that the difference between the rolling temperature in the first pass and the rolling temperature in the last pass in the rolling in a state where the central portion of the plate thickness is in the austenite non-recrystallized region is within 40 ° C. Here, the rolling temperature refers to the temperature at the center of the plate thickness of the steel just before rolling. The temperature at the center of the plate thickness is obtained by simulation calculation or the like from the plate thickness, surface temperature, thermal history, and the like. For example, the temperature at the center of the plate thickness of the steel sheet is obtained by calculating the temperature distribution in the plate thickness direction using the difference method.
　上記のオーステナイト再結晶温度域およびオーステナイト未再結晶温度域を合わせた合計の累積圧下率は６５％以上とすることが好ましい。全体の圧下率が小さいと、組織の圧下が十分でなく、靭性および強度が目的の値を達成することが出来ない。全体の累積圧下率を６５％以上とすることにより、組織に対して十分な圧下量を確保することができ、靭性および集積度が目的の値を達成することができる。 It is preferable that the total cumulative rolling reduction of the austenite recrystallization temperature range and the austenite non-recrystallization temperature range be 65% or more. When the overall rolling reduction is small, the rolling of the structure is not sufficient, and the toughness and strength cannot achieve the target values. By setting the total cumulative reduction ratio to 65% or more, a sufficient amount of reduction can be ensured for the structure, and the toughness and the degree of accumulation can achieve the target values.
　オーステナイト再結晶温度域、および、オーステナイト未再結晶温度域は、当該成分組成を有する鋼に、条件を変化させた熱・加工履歴を与える予備的実験を行うことにより、把握することができる。 The austenite recrystallization temperature range and the austenite non-recrystallization temperature range can be grasped by conducting a preliminary experiment in which the steel having the component composition is given a heat / working history with varying conditions.
　なお、熱間圧延の終了温度は特に限定されるものではない。圧延能率の観点からは、オーステナイト未再結晶温度域において終了させることが好ましい。 Note that the end temperature of hot rolling is not particularly limited. From the viewpoint of rolling efficiency, it is preferable to terminate in the austenite non-recrystallization temperature range.
　圧延が終了した鋼板は、４．０℃／ｓ以上の冷却速度にて４５０℃以下まで冷却することが好ましい。冷却速度を４．０℃／ｓ以上とすることにより、組織が粗大化することなく、また、フェライト変態を抑制することにより、細粒のベイナイト組織が得られ、目標とする優れた靱性や集積度を得ることができる。冷却速度が４．０℃／ｓ未満では、各板厚位置において、組織の粗大化やフェライト変態が進むため、所望の組織が得られないばかりか、鋼板の強度も低下する。 It is preferable that the rolled steel sheet is cooled to 450 ° C. or lower at a cooling rate of 4.0 ° C./s or higher. By setting the cooling rate to 4.0 ° C./s or more, the microstructure does not become coarse, and by suppressing the ferrite transformation, a fine-grained bainite structure can be obtained, and the excellent excellent toughness and accumulation can be achieved. You can get a degree. When the cooling rate is less than 4.0 ° C./s, the coarsening of the structure and ferrite transformation proceed at each plate thickness position, so that a desired structure cannot be obtained and the strength of the steel sheet also decreases.
　冷却停止温度を４５０℃以下とすることにより、ベイナイト変態を十分に進行されることができ、所望の靭性や集積度を得ることができる。冷却停止温度が４５０℃より高いと、ベイナイト変態が十分には進行せず、フェライトやパーライトなどの組織も生成し、本発明が目的とするベイナイト主体の組織が得られない。なお、これら冷却速度や冷却停止温度は、鋼板の板厚中央部の温度とする。板厚中央部の温度は、板厚、表面温度および冷却条件等から、シミュレーション計算等により求められる。例えば、差分法を用い、板厚方向の温度分布を計算することにより、鋼板の板厚中央部の温度が求められる。 When the cooling stop temperature is set to 450 ° C. or lower, the bainite transformation can be sufficiently advanced, and desired toughness and integration degree can be obtained. If the cooling stop temperature is higher than 450 ° C., the bainite transformation does not proceed sufficiently, and a structure such as ferrite or pearlite is also produced, and the bainite-based structure intended by the present invention cannot be obtained. In addition, let these cooling rate and cooling stop temperature be the temperature of the plate | board thickness center part of a steel plate. The temperature at the center of the plate thickness is obtained by simulation calculation or the like from the plate thickness, surface temperature, cooling conditions, and the like. For example, the temperature at the center of the plate thickness of the steel sheet is obtained by calculating the temperature distribution in the plate thickness direction using the difference method.
　冷却が終了した鋼板について、焼戻し処理を実施することも可能である。焼き戻しを実施することにより、鋼板の靭性をさらに向上させることができる。焼戻し温度は、鋼板平均温度でＡＣ１点以下として、焼戻し処理を実施することにより、圧延・冷却で得られた所望の組織を損なわないようにすることができる。本発明ではＡＣ１点（℃）を下式で求める。
ＡＣ１点＝７５１－２６．６Ｃ＋１７．６Ｓｉ－１１．６Ｍｎ－１６９Ａｌ－２３Ｃｕ－２３Ｎｉ＋２４．１Ｃｒ＋２２．５Ｍｏ＋２３３Ｎｂ－３９．７Ｖ－５．７Ｔｉ－８９５Ｂ
上記式において、各元素記号は鋼中含有量（質量％）であり、含有しない場合は０とする。 It is also possible to perform a tempering process on the steel plate that has been cooled. By performing tempering, the toughness of the steel sheet can be further improved. Tempering temperature as follows C1 points A steel plate average temperature, by carrying out the tempering treatment, it is possible not impair the desired tissue obtained by rolling and cooling. In the present invention, the AC1 point (° C.) is obtained by the following equation.
A C1 point = 751-26.6C + 17.6Si-11.6Mn-169Al-23Cu-23Ni + 24.1Cr + 22.5Mo + 233Nb-39.7V-5.7Ti-895B
In the above formula, each element symbol is the content (% by mass) in steel, and 0 if not contained.
　鋼板の平均温度も、板厚中央部の温度と同様、板厚、表面温度および冷却条件等から、シミュレーション計算等により求められる。 The average temperature of the steel sheet can also be obtained by simulation calculation or the like from the sheet thickness, surface temperature, cooling conditions, etc., similarly to the temperature at the center of the sheet thickness.
　表１に示す各組成の溶鋼（鋼記号Ａ～Ｑ）を転炉で溶製し、連続鋳造法で鋼素材（スラブ２５０ｍｍ厚または３００ｍｍ厚）とし、板厚５５～１００ｍｍに熱間圧延後、冷却を行い、Ｎｏ．１～２７の供試鋼を得た。一部については、冷却後に焼戻しも実施した。表２に熱間圧延条件と冷却条件を示す。 Molten steel (steel symbols A to Q) of each composition shown in Table 1 was melted in a converter and made into a steel material (slab 250 mm thickness or 300 mm thickness) by a continuous casting method. After hot rolling to a plate thickness of 55 to 100 mm, Cooling is performed, and no. Sample steels of 1 to 27 were obtained. Some were tempered after cooling. Table 2 shows hot rolling conditions and cooling conditions.
　得られた厚鋼板について、板厚の１／４部よりφ１４ｍｍのＪＩＳ１４Ａ号試験片を試験片の長手方向が圧延方向と直角になるように採取し、引張試験を行い、降伏強度（ＹＳ）および引張強さ（ＴＳ）を測定した。 About the obtained thick steel plate, a JIS14A test piece having a diameter of 14 mm was collected from 1/4 part of the plate thickness so that the longitudinal direction of the test piece was perpendicular to the rolling direction, a tensile test was performed, and the yield strength (YS) and Tensile strength (TS) was measured.
　また、靭性値を評価するため、板厚表層部および板厚中央部（以下、板厚中央部を１／２ｔ部と記す場合がある。）よりＪＩＳ４号衝撃試験片を試験片の長手方向が圧延方向と平行になるように採取し、シャルピー衝撃試験を行って破面遷移温度（ｖＴｒｓ）をそれぞれ求めた。ここで、表層部の衝撃試験片は、最も表面に近い面を鋼板表面から１ｍｍの深さにするものとする。 Further, in order to evaluate the toughness value, the longitudinal direction of the test piece was measured in accordance with the JIS No. 4 impact test piece from the plate thickness surface layer portion and the plate thickness central portion (hereinafter, the plate thickness central portion may be referred to as a 1/2 t portion). Samples were taken so as to be parallel to the rolling direction, and Charpy impact tests were performed to determine fracture surface transition temperatures (vTrs). Here, the impact test piece of the surface layer part is assumed to have a surface closest to the surface at a depth of 1 mm from the steel sheet surface.
　得られた厚鋼板の圧延長手方向と平行な板厚断面を鏡面研磨したあと、エッチングにより現出させた金属組織を光学顕微鏡によって観察した。 After the plate thickness section parallel to the rolling longitudinal direction of the obtained thick steel plate was mirror-polished, the metal structure revealed by etching was observed with an optical microscope.
　次に、脆性亀裂伝播停止特性を評価するために、標準ＥＳＳＯ試験（温度勾配型ＥＳＳＯ試験）を行い、－１０℃におけるＫｃａ値（Ｋｃａ（－１０℃））を求めた。 Next, in order to evaluate the brittle crack propagation stop property, a standard ESSO test (temperature gradient type ESSO test) was performed to obtain a Kca value (Kca (−10 ° C.)) at −10 ° C.
　さらに、板厚中央部におけるＲＤ／／（１１０）面の集積度を次のようにして求めた。まず、板厚中央部から板厚１ｍｍのサンプルを採取し、板面に平行な面を機械研磨・電解研磨することにより、Ｘ線回折用の試験片を用意した。この試験片を用いて、Ｍｏ線源を用いてＸ線回折装置を用いて、Ｘ線回折測定を実施し、（２００）、（１１０）および（２１１）正極点図を求め、得られた正極点図から３次元結晶方位分布密度関数をＢｕｎｇｅ法で計算して求める。次に得られた３次元結晶方位分布密度関数からφ２＝０～９０°まで、Ｂｕｎｇｅ表記で５°間隔で合計１９枚の断面図において、圧延方向に対して（１１０）面が平行となる方位の3次元結晶方位分布密度関数の値を積算して積算値を求め、この積算値を前記積算した方位の個数で割った値を、ＲＤ／／（１１０）面の集積度とした。 Furthermore, the degree of integration of the RD // (110) plane at the center of the plate thickness was determined as follows. First, a sample having a plate thickness of 1 mm was collected from the central portion of the plate thickness, and a test piece for X-ray diffraction was prepared by mechanically polishing and electrolytic polishing a surface parallel to the plate surface. Using this test piece, X-ray diffraction measurement was performed using an X-ray diffractometer using a Mo ray source, and (200), (110) and (211) positive electrode dot diagrams were obtained, and the obtained positive electrode A three-dimensional crystal orientation distribution density function is calculated from the dot diagram by the Bunge method. Next, from the obtained three-dimensional crystal orientation distribution density function, φ2 = 0 to 90 °, orientation in which the (110) plane is parallel to the rolling direction in a total of 19 cross-sectional views at 5 ° intervals in Bunge notation The integrated value of the three-dimensional crystal orientation distribution density function is integrated to obtain an integrated value, and the integrated value of the RD // (110) plane is obtained by dividing the integrated value by the number of the integrated orientations.
　大入熱溶接特性を評価するために、供試鋼板に開先加工を施し（開先角度２０°）、市販の低温用鋼用エレクトロガスアーク溶接用ワイヤを使用してエレクトロガス溶接で入熱３００～７５０ｋＪ／ｃｍで溶接継手を作製し、ＨＡＺ靭性として、ボンド部の靱性を２ｍｍＶノッチシャルピー試験により評価した。試験は、－２０℃でのシャルピー吸収エネルギーのｖＥ－２０（３本平均値）で行った。 In order to evaluate the large heat input welding characteristics, the test steel sheet was subjected to groove processing (groove angle 20 °), and heat input 300 was performed by electrogas welding using a commercially available wire for electrogas arc welding for low temperature steel. A welded joint was prepared at ˜750 kJ / cm, and the toughness of the bond portion was evaluated by a 2 mmV notch Charpy test as HAZ toughness. The test was performed with vE -20 (average value of three) of Charpy absorbed energy at -20 ° C.
　表３にこれらの試験結果を示す。本発明の範囲内にある供試鋼板（製造Ｎｏ．１～１１）は、Ｋｃａ（－１０℃）が６０００Ｎ／ｍｍ３／２以上の優れた脆性亀裂伝播停止性能を示した。また、大入熱溶接継手のボンド部の吸収エネルギー：ｖＥ－２０≧８８Ｊとなり、優れた値を示した。また、表層部および板厚中央部のシャルピー靭性値（破面遷移温度）、および、ＲＤ／／（１１０）集積度が（２）式を満たしている供試鋼板（製造番号２～１１）においては、（２）式を満たしていない供試鋼板（製造番号１）と比較して、高いＫｃａ（－１０℃）の値が得られた。なお、これらの供試鋼板（製造Ｎｏ．１～１１）の金属組織は、いずれもベイナイト主体であった。 Table 3 shows the results of these tests. The test steel plates (production Nos. 1 to 11) within the scope of the present invention exhibited excellent brittle crack propagation stopping performance with a Kca (−10 ° C.) of 6000 N / mm 3/2 or more. Further, the absorbed energy of the bond portion of the high heat input welded joint was vE-20 ≧ 88 J, which was an excellent value. Further, in the test steel plates (manufacturing numbers 2 to 11) in which the Charpy toughness value (fracture surface transition temperature) of the surface layer portion and the center portion of the plate thickness and the RD // (110) accumulation degree satisfy the formula (2) Compared with the test steel plate (Production No. 1) not satisfying the formula (2), a higher Kca (−10 ° C.) value was obtained. Note that the metal structures of these test steel sheets (production Nos. 1 to 11) were mainly bainite.
　一方、鋼板の成分は本発明の好ましい範囲であるものの、鋼板の製造条件における加熱、圧延条件が本発明の好ましい範囲を外れる鋼板（製造Ｎｏ．２０～２７）は、Ｋｃａ（－１０℃）の値は、６０００Ｎ／ｍｍ３／２には達しなかった。鋼板の成分が本発明の条件を満たさない供試鋼板（製造Ｎｏ．１２～１９）については、大入熱溶接継手の吸収エネルギー：ｖＥ－２０が２２Ｊ以下となり、本発明例と比較して劣った。 On the other hand, although the steel plate component is within the preferred range of the present invention, the steel plate (production Nos. 20 to 27) whose heating and rolling conditions are outside the preferred range of the present invention is Kca (−10 ° C.). The value did not reach 6000 N / mm 3/2 . For the test steel sheets (production Nos. 12 to 19) in which the components of the steel sheet do not satisfy the conditions of the present invention, the absorbed energy of the large heat input welded joint: vE- 20 is 22 J or less, which is inferior to the present invention example. It was.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
１　　標準ＥＳＳＯ試験片
２　　ノッチ
３　　亀裂
３ａ　分岐
４　　先端形状
５　　母材 1 Standard ESSO Test Piece 2 Notch 3 Crack 3a Branch 4 Tip Shape 5 Base Material

Claims (5)

  1. 　鋼組成が、質量％で、Ｃ：０．０３～０．１５％、Ｓｉ：０．０１～０．５％、Ｍｎ：１．４０～２．５０％、Ａｌ：０．００５～０．０８％、Ｐ：０．０３％以下、Ｓ：０．０００５～０．００３０％、Ｎ：０．００３６～０．００７０％、Ｔｉ：０．００４～０．０３０％、Ｃａ：０．０００５～０．００３０％を含有し、且つ、Ｃａ、Ｓ、Ｏの各含有量が、下記（１）式を満足し、残部がＦｅおよび不可避的不純物で、金属組織がベイナイトを主体とし、板厚中央部におけるＲＤ／／（１１０）面の集積度が１．５～４．０の集合組織を有し、かつ表層部および板厚中央部におけるシャルピー破面遷移温度ｖＴｒｓが－４０℃以下であることを特徴とする脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。
    ０．３０≦（Ｃａ－（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５／Ｓ≦０．８０・・・（１）
    ただし、式（１）において、Ｃａ、Ｏ、Ｓは含有量（質量％）とする。     Steel composition is mass%, C: 0.03-0.15%, Si: 0.01-0.5%, Mn: 1.40-2.50%, Al: 0.005-0.08 %, P: 0.03% or less, S: 0.0005 to 0.0030%, N: 0.0036 to 0.0070%, Ti: 0.004 to 0.030%, Ca: 0.0005 to 0 .0030%, and each content of Ca, S, and O satisfies the following formula (1), the balance is Fe and inevitable impurities, the metal structure is mainly bainite, and the central portion of the plate thickness The RD // (110) plane has a texture of 1.5 to 4.0, and the Charpy fracture surface transition temperature vTrs in the surface layer portion and the thickness center portion is −40 ° C. or lower. High strength thick steel plate for high heat input welding with excellent brittle crack propagation stopping characteristics.
    0.30 ≦ (Ca− (0.18 + 130 × Ca) × O) /1.25/S≦0.80 (1)
    However, in Formula (1), Ca, O, and S are made into content (mass%).
  2. 　鋼組成が、更に、質量％で、Ｎｂ：０．０５％以下、Ｃｕ：１．０％以下、Ｎｉ：１．０％以下、Ｃｒ：０．５％以下、Ｍｏ：０．５％以下、Ｖ：０．２％以下、Ｂ：０．００３％以下、ＲＥＭ：０．０１％以下の１種または２種以上を含有することを特徴とする請求項１に記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。 The steel composition is further mass%, Nb: 0.05% or less, Cu: 1.0% or less, Ni: 1.0% or less, Cr: 0.5% or less, Mo: 0.5% or less, V: 0.2% or less, B: 0.003% or less, REM: 0.01% or less, or one or two or more of them are contained. Excellent high strength thick steel plate for high heat input welding.
  3. 　板厚中央部のシャルピー破面遷移温度およびＲＤ／／（１１０）面の集積度が、下記（２）式を満たすことを特徴とする請求項１または２記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板。
    ｖＴｒｓ（１／２ｔ）－１２×ＩＲＤ／／（１１０）［１／２ｔ］≦－７０・・・（２）
    ただし、式（２）において、
    ｖＴｒｓ（１／２ｔ）：板厚中央部（１／２ｔ）のシャルピー破面遷移温度（℃）
    ＩＲＤ／／（１１０）［１／２ｔ］：板厚中央部（１／２ｔ）のＲＤ／／（１１０）面の集積度
    とする。なお、ｔは板厚（ｍｍ）である。     The Charpy fracture surface transition temperature in the central portion of the plate thickness and the degree of integration of the RD // (110) surface satisfy the following formula (2): Excellent brittle crack propagation stop characteristics according to claim 1 or 2 High strength thick steel plate for high heat input welding.
    vTrs (1 / 2t) −12 × I RD // (110) [1 / 2t] ≦ −70 (2)
    However, in Formula (2),
    vTrs (1 / 2t) : Charpy fracture surface transition temperature (° C.) at the center of the plate thickness (1 / 2t)
    I RD // (110) [1 / 2t] : The degree of integration of the RD // (110) plane in the central portion (1 / 2t) of the plate thickness. Here, t is a plate thickness (mm).
  4. 　請求項１または２に記載の組成を有する鋼素材を、１０００～１２００℃の温度に加熱し、オーステナイト再結晶温度域およびオーステナイト未再結晶温度域における累積圧下率の合計が６５％以上の圧延を実施し、このとき、板厚中央部がオーステナイト再結晶温度域にある状態においては累積圧下率が２０％以上の圧延を行い、次いで、板厚中央部がオーステナイト未再結晶温度域にある状態においては、累積圧下率が４０～７０％とする圧延を行い、かつ、前記板厚中央部がオーステナイト未再結晶温度域にある状態における圧延のうち最初のパスの圧延温度と最後のパスの圧延温度との差が４０℃以内であり、その後、４．０℃／ｓ以上の冷却速度にて４５０℃以下まで冷却することを特徴とする脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板の製造方法。 The steel material having the composition according to claim 1 or 2 is heated to a temperature of 1000 to 1200 ° C., and rolling with a total cumulative rolling reduction of 65% or more in the austenite recrystallization temperature range and the austenite non-recrystallization temperature range is performed. At this time, in the state where the plate thickness central portion is in the austenite recrystallization temperature range, the rolling reduction is performed at a cumulative reduction ratio of 20% or more, and then in the state where the plate thickness central portion is in the austenite non-recrystallization temperature range. Is the rolling temperature of the first pass and the rolling temperature of the last pass of the rolling in a state where the rolling reduction is 40 to 70% and the central part of the plate thickness is in the austenite non-recrystallization temperature range. Is excellent in brittle crack propagation stop characteristics characterized by cooling to 450 ° C. or lower at a cooling rate of 4.0 ° C./s or higher. Method of producing a high strength steel plate for hot welding.
  5. 　４５０℃以下に加速冷却した後、さらに、Ａｃ１点以下の温度に焼戻す工程を有することを特徴とする請求項４に記載の脆性亀裂伝播停止特性に優れた大入熱溶接用高強度厚鋼板の製造方法。 The high-strength thickness for high heat input welding having excellent brittle crack propagation stopping characteristics according to claim 4, further comprising a step of tempering to a temperature not higher than Ac1 point after accelerated cooling to 450 ° C or lower. A method of manufacturing a steel sheet.
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Cited By (4)

* Cited by examiner, † Cited by third party
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WO2017145651A1 (en) * 2016-02-24 2017-08-31 Ｊｆｅスチール株式会社 High strength ultra-thick steel plate having excellent brittle crack propagation stopping characteristics and manufaturing method of same
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WO2023219146A1 (en) * 2022-05-12 2023-11-16 Ｊｆｅスチール株式会社 Steel sheet and method for manufacturing same

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JP6048626B1 (en) * 2015-01-16 2016-12-21 Ｊｆｅスチール株式会社 Thick, high toughness, high strength steel plate and method for producing the same
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Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2007302993A (en) * 2006-04-13 2007-11-22 Nippon Steel Corp High-strength steel plate with superior arrestability
JP2008169468A (en) * 2006-12-14 2008-07-24 Nippon Steel Corp High-strength thick steel plate having excellent brittle crack propagation-stopping performance
JP2008169467A (en) * 2006-12-14 2008-07-24 Nippon Steel Corp High-strength thick steel plate having excellent brittle crack propagation-stopping performance, and method for producing the same
JP2008248382A (en) * 2007-03-05 2008-10-16 Nippon Steel Corp Thick high-strength steel plate having superior brittle-crack propagation-stopping characteristic and manufacturing method therefor
JP2009019244A (en) * 2007-07-12 2009-01-29 Kobe Steel Ltd Thick steel plate for high heat input welding having excellent brittle fracture propagation stop characteristic
WO2009072559A1 (en) * 2007-12-06 2009-06-11 Nippon Steel Corporation Process for producing thick high-strength steel plate excellent in brittle fracture arrestability and toughness of zone affected by heat in large-heat-input welding and thick high-strength steel plate excellent in brittle fracture arrestability and toughness of zone affected by heat in large-heat-input welding

Family Cites Families (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP4901623B2 (en) 2007-07-20 2012-03-21 新日本製鐵株式会社 High-strength steel sheet with excellent punching hole expandability and manufacturing method thereof
JP5434145B2 (en) * 2009-03-04 2014-03-05 Ｊｆｅスチール株式会社 Structural high-strength thick steel plate with excellent brittle crack propagation stopping characteristics and method for producing the same
JP5477578B2 (en) 2010-04-01 2014-04-23 新日鐵住金株式会社 Thick high-strength steel sheet excellent in brittle crack propagation stopping characteristics and method for producing the same
JP5304925B2 (en) * 2011-12-27 2013-10-02 Ｊｆｅスチール株式会社 Structural high-strength thick steel plate with excellent brittle crack propagation stopping characteristics and method for producing the same

Patent Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2007302993A (en) * 2006-04-13 2007-11-22 Nippon Steel Corp High-strength steel plate with superior arrestability
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CN105102650B (en) 2017-10-24
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KR101732997B1 (en) 2017-05-08
CN105102650A (en) 2015-11-25

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