EP4520845A1 - Hochfestes stahlblech für ein wasserstofftransportstahlrohr, herstellungsverfahren dafür und wasserstofftransportstahlrohr - Google Patents

Hochfestes stahlblech für ein wasserstofftransportstahlrohr, herstellungsverfahren dafür und wasserstofftransportstahlrohr Download PDF

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EP4520845A1
EP4520845A1 EP23839276.5A EP23839276A EP4520845A1 EP 4520845 A1 EP4520845 A1 EP 4520845A1 EP 23839276 A EP23839276 A EP 23839276A EP 4520845 A1 EP4520845 A1 EP 4520845A1
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Prior art keywords
steel plate
less
temperature
steel
content
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English (en)
French (fr)
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EP4520845A4 (de
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Daichi IZUMI
Yoshihiro Nishihara
Hiroshi Okano
Junji Shimamura
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JFE Steel Corp
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JFE Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • C21D9/085—Cooling or quenching
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • 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
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/20—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
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—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

Definitions

  • the present invention relates to high strength steel plates for hydrogen transport steel pipes, and more particularly to a high strength steel plate for a hydrogen transport steel pipe suitable for use in a line pipe for transporting high-pressure hydrogen gas, and a method for manufacturing the high strength steel plate.
  • the present invention also relates to a steel pipe for transporting hydrogen, the steel pipe using the above-mentioned high strength steel plate for transporting hydrogen.
  • Line pipes are typically produced by forming steel plates, which are produced by plate mills or hot rolling mills, into steel pipes using, for example, UOE forming, press bend forming, or roll forming.
  • Line pipes used for transporting high-pressure hydrogen gas are required to have hydrogen embrittlement resistance in addition to strength, toughness, weldability, and so forth.
  • fatigue crack growth resistance in a high-pressure hydrogen gas environment is required to extend their service life.
  • hydrogen-induced stress cracking resistance (HISC resistance) in a high-pressure hydrogen gas environment is required.
  • HISC resistance hydrogen-induced stress cracking resistance
  • the hydrogen pressure is about 15 MPa
  • low-alloy steel having a sufficient wall thickness is used.
  • austenitic stainless steel such as SUS316L, which is less susceptible to hydrogen embrittlement than low-alloy steel, is used.
  • Austenitic stainless steel has low strength in addition to the high cost of the steel material. Therefore, if austenitic stainless steel is designed to withstand high hydrogen pressure, the wall thickness of the austenitic stainless steel will be increased, and the hydrogen transport line pipe itself will also be more expensive. For this reason, there has been a demand for a steel material for hydrogen transport line pipes that is lower in cost and can withstand a high-pressure hydrogen gas environment.
  • Patent Literature 1 discloses an austenitic steel material having a high Mn content.
  • Patent Literature 1 According to the technique described in Patent Literature 1, it is possible to provide a steel material that is lower in cost than austenitic stainless steel such as SUS316L.
  • the steel material described in Patent Literature 1 is an austenitic alloy and thus is more expensive than a typical low-alloy steel.
  • HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen gas environment are not taken into consideration.
  • the inventors have repeatedly conducted numerous experiments and studies on the chemical compositions, microstructures, and production conditions of steel materials in order to ensure HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen gas environment, and have found the following. That is, letting the standard deviation of a Vickers hardness at 0.25 mm below a surface of a steel plate be denoted as ⁇ , the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3 ⁇ is controlled to 225 HV or less, and the top 20% grain size in the microstructure at the center of the plate thickness is 30 ⁇ m or less. This improves HISC resistance and fatigue crack growth resistance. To achieve such a steel microstructure, it is necessary to strictly control the rolling and cooling conditions. The inventors have succeeded in finding those conditions. The present invention has been made based on these findings.
  • the gist and configuration of the present invention are as described below.
  • the high strength steel plate of the present invention for the hydrogen transport steel pipe and the steel pipe for transporting hydrogen using the high strength steel plate for the hydrogen transport steel pipe have excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment.
  • the steel pipe has excellent HISC resistance in a high-pressure hydrogen environment, even in a region including the weld of the steel pipe.
  • According to the method of the present invention for manufacturing a high strength steel plate for a hydrogen transport steel pipe it is possible to manufacture a high strength steel plate for a hydrogen transport steel pipe, the high strength steel plate having excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment.
  • Fig. 1 is a schematic view illustrating a method for collecting a test specimen for evaluating HISC resistance in examples.
  • a high strength steel plate of the present invention for a hydrogen transport steel pipe will be specifically described below.
  • the high strength steel plate of the present invention for a hydrogen transport steel pipe is also referred to simply as a "high strength steel plate”.
  • the C content effectively contributes to improving strength.
  • the C content is 0.030% or more.
  • the C content is preferably 0.035% or more.
  • the C content is more than 0.060%, the hardness increases during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the C content is 0.060% or less.
  • the C content is preferably 0.050% or less.
  • Si is added for deoxidation.
  • a Si content of less than 0.01% results in an insufficient deoxidation effect. For this reason, the Si content is 0.01% or more.
  • the Si content is preferably 0.05% or more.
  • a Si content of more than 0.50% results in a deterioration in weldability. For this reason, the Si content is 0.50% or less.
  • the Si content is preferably 0.45% or less.
  • Mn effectively contributes to improving strength. However, when the Mn content is less than 0.80%, this effect is not sufficiently provided. For this reason, the Mn content is 0.80% or more.
  • the Mn content is preferably 1.00% or more.
  • the Mn content is more preferably 1.20% or more.
  • a Mn content of more than 1.80% results in an increase in hardness during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the Mn content is 1.80% or less.
  • the Mn content is preferably 1.70% or less.
  • the Mn content is more preferably 1.60% or less.
  • P is an incidental impurity element that increases the hardness and thereby deteriorates the HISC resistance.
  • the upper limit of the P content is 0.015%.
  • the P content is preferably 0.008% or less.
  • a lower P content is more preferred.
  • excessive dephosphorization leads to an increase in refining cost. From the viewpoint of the refining cost, the P content is preferably 0.001% or more.
  • S is an incidental impurity element that forms MnS inclusions in steel to deteriorate low-temperature toughness. For this reason, a lower S content is more preferred. However, a S content of up to 0.0015% is acceptable. For this reason, the S content is 0.015% or less. The S content is preferably 0.0010% or less. A lower S content is more preferred. However, excessive desulfurization leads to an increase in refining cost. From the viewpoint of the refining cost, the S content is preferably 0.0002% or more.
  • Al is added as a deoxidizing agent.
  • the Al content is less than 0.010%, the effect is not sufficiently provided.
  • the Al content is 0.010% or more.
  • the Al content is preferably 0.015% or more.
  • the Al content is more preferably 0.025% or more.
  • An Al content of more than 0.080% results in clogging of a submerged nozzle with alumina during continuous casting. For this reason, the Al content is 0.080% or less.
  • the Al content is preferably 0.070% or less.
  • the Al content is more preferably 0.040% or less.
  • Cr is an effective element for achieving sufficient strength even in steel having a low C content.
  • the Cr content is preferably 0.10% or more.
  • the Cr content is more preferably 0.15% or more.
  • a Cr content of more than 0.50% results in excessive hardenability to increase the hardness during accelerated cooling, thereby deteriorating the HISC resistance.
  • the Cr content is 0.50% or less.
  • the Cr content is preferably 0.45% or less.
  • the Cr content is more preferably 0.35% or less.
  • Nb When Nb is present as solute Nb, Nb extends the non-recrystallization temperature range during hot rolling and contributes to refinement of crystal grain size. When the Nb content is less than 0.005%, this effect is not sufficiently provided. For this reason, the Nb content is 0.005% or more.
  • the Nb content is preferably 0.010% or more.
  • the Nb content is more preferably 0.025% or more.
  • a Nb content of more than 0.080% results in crystallization of coarse carbides during solidification, deteriorating hydrogen induced cracking resistance. For this reason, the Nb content is 0.080% or less.
  • the Nb content is preferably 0.060% or less.
  • the Nb content is more preferably 0.055% or less.
  • Ti has the effect of pinning austenite grains in the form of TiN during heating, thereby inhibiting the growth of the grains.
  • Ti content is less than 0.005%, TiN is not sufficiently formed.
  • the Ti content is 0.005% or more.
  • the Ti content is preferably 0.008% or more.
  • the Ti content is more than 0.020%, the formed TiN coarsens, resulting in a failure to obtain sufficient toughness of the weld heat affected zone. For this reason, the Ti content is 0.020% or less.
  • the Ti content is preferably 0.017% or less.
  • the Ti content is more preferably 0.015% or less.
  • N effectively contributes to improving strength.
  • the N content is 0.0020% or more.
  • the N content is preferably 0.0025% or more.
  • the N content is more preferably 0.0030% or more.
  • a N content of more than 0.0080% results in an increase in hardness during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the N content is 0.0080% or less.
  • the N content is preferably 0.0070% or less.
  • the N content is more preferably 0.0050% or less.
  • Ca is an element effective in improving hydrogen induced cracking resistance by shape control of sulfide-based inclusions.
  • the Ca content is preferably 0.0008% or more.
  • the Ca content is more preferably 0.0015% or more.
  • a Ca content of more than 0.0050% results in saturation of the above-mentioned effect and a decrease in the cleanliness of the steel, deteriorating the hydrogen induced cracking resistance. For this reason, the Ca content is 0.0050% or less.
  • the Ca content is preferably 0.0045% or less.
  • the Ca content is more preferably 0.0035% or less.
  • components other than those mentioned above can be Fe and incidental impurities.
  • the chemical composition of the high strength steel plate of the present invention can further contain, in addition to the above-mentioned components, one or more selected from Cu, Ni, Mo, V, Zr, Mg, and REM within the following ranges.
  • the Cu is an element effective in improving low-temperature toughness and increasing strength. To provide these effects, the Cu content is preferably 0.05% or more. The Cu content is more preferably 0.10% or more. When the Cu content is more than 0.50%, defects on the steel plate surface are likely to occur. For this reason, when Cu is contained, the Cu content is 0.50% or less. The Cu content is preferably 0.45% or less.
  • Ni is an element effective in improving low-temperature toughness and increasing strength. To provide these effects, the Ni content is preferably 0.05% or more. The Ni content is more preferably 0.10% or more. Ni is an expensive element. For this reason, when Ni is contained, the Ni content is 0.50% or less. The Ni content is preferably 0.45% or less.
  • Mo is an element in improving low-temperature toughness and increasing strength. To provide these effects, the Mo content is preferably 0.05% or more. Mo is an expensive element. For this reason, when Mo is contained, the Mo content is 0.50% or less. The Mo content is preferably 0.45% or less.
  • V is an element that can be added to the steel plate to enhance its strength and low-temperature toughness.
  • the V content is less than 0.005%, the effect is not sufficiently provided.
  • the V content is preferably 0.005% or more.
  • a V content of more than 0.1% results in a deterioration in the toughness of the weld.
  • the V content is preferably 0.1% or less.
  • the V content is more preferably 0.050% or less, still more preferably 0.010% or less.
  • Zr, Mg, and a rare earth metal (REM) are elements that can be optionally added to enhance fatigue crack growth resistance through the refinement of crystal grains and to enhance crack resistance through control of inclusion properties.
  • the amount of each element contained is less than 0.0005%, the effects are not sufficiently provided. For this reason, when these elements are contained, the amount of each element contained is preferably 0.0005% or more. When the amount of each element contained is more than 0.02%, the effects are saturated. For this reason, when Zr, Mg, and REM are contained, the amount of each element contained is preferably 0.02% or less.
  • the amount of each of the above elements contained is more preferably 0.0050% or less, still more preferably 0.0030% or less.
  • REM is a general term for Sc, Y, and 15 elements from lanthanum (La) of atomic number 57 to lutetium (Lu) of atomic number 71. The amount of REM contained refers to the total amount of these elements.
  • the balance other than the above elements is Fe and incidental impurities.
  • other trace elements may be contained as long as they do not impair the effects of the present invention.
  • O is an element that is incidentally contained in steel, and is acceptable in the present invention as long as the O content is 0.0050% or less, preferably 0.0040% or less.
  • the "Vickers hardness (HV 0.5) at 0.25 mm below the surface of the steel plate” refers to the Vickers hardness (HV 0.5) measured at 100 equally spaced points in the plate width direction from each of the leading end and the trailing end in the rolling direction of the steel plate, the 100 equally spaced points being located at 0.25 mm below the surface of the steel plate (at a depth of 0.25 mm from the surface of the steel plate toward the center of the plate thickness).
  • the leading end of the steel plate in the rolling direction is a position 1 m downstream of the leading edge of the steel plate in the rolling direction.
  • the trailing end of the steel plate in the rolling direction is a position 1 m upstream of the trailing edge of the steel plate in the rolling direction.
  • the measurement was performed in a region excluding non-steady portions near the ends in the plate width direction.
  • the reason for measuring the hardness of the steel plate at 0.5 kgf instead of the commonly used 10 kgf is that measurement at 0.5 kgf results in a smaller indentation, making it possible to obtain hardness information at a position closer to the surface and hardness information that is more sensitive to the microstructure.
  • Measuring the Vickers hardness with a test load of less than 0.5 kgf results in a too small indentation size and thus an increase in measurement variability, which is not preferred.
  • the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3 ⁇ is preferably 220 HV or less.
  • the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate +3 ⁇ is 200 HV or more.
  • the top 20% grain size is preferably 25 ⁇ m or less. As an example, the top 20% grain size is 15 ⁇ m or more.
  • the top 20% grain size is a grain size corresponding to the 20% position from the largest crystal grain size when the crystal grain sizes are arranged in descending order in the crystal grain size distribution.
  • the measurement range of the crystal grain size was 1 mm ⁇ 1 mm at the center position of the plate thickness. More specifically, the crystal grain size was determined as follows: The microstructure at the center of the plate thickness was analyzed by an electron backscatter diffraction (EBSD) method. A boundary having an orientation difference of 15° or more was determined to be a crystal grain boundary. The equivalent circular diameter was calculated as the crystal grain size from the area of each crystal grain.
  • EBSD electron backscatter diffraction
  • a frequency distribution table is prepared for all crystal grains to be measured, and the crystal grain size corresponding to 20% of the cumulative relative frequency from the largest crystal grain size calculated is referred to as the "top 20% grain size".
  • the fatigue crack growth rate is less than 2.0 ⁇ 10 -2 (mm/cycle) when the stress intensity factor range ⁇ K is 45 (MPa ⁇ m 1/2 ).
  • the fatigue crack growth rate is preferably 1.5 ⁇ 10 -2 (mm/cycle) or less.
  • a lower fatigue crack growth rate is more preferred.
  • the fatigue crack growth rate is 1.0 ⁇ 10 -2 (mm/cycle) or more.
  • the high strength steel plate of the present invention is intended mainly for steel plates for steel pipes having a strength of API 5L grade X65 or higher, and thus has a tensile strength of 535 MPa or higher.
  • the upper limit of the tensile strength of the high strength steel plate of the present invention is not particularly limited, but as an example, the tensile strength of the high strength steel plate of the present invention is 760 MPa or less.
  • the high strength steel plate of the present invention may have a tensile strength of 600 MPa or less.
  • the plate thickness of the high strength steel plate of the present invention is not particularly limited, but is preferably 12 mm or more.
  • the plate thickness of the high strength steel plate of the present invention is not particularly limited, but is preferably 39 mm or less.
  • a manufacturing method and manufacturing conditions for manufacturing the above-mentioned high strength steel plate are specifically described below.
  • the method for manufacturing the high strength steel plate of the present invention includes heating a steel slab (slab) having the above-mentioned chemical composition, then subjecting the steel slab to hot rolling to form a steel plate (hot rolling step), and thereafter cooling the steel plate under predetermined conditions (cooling step).
  • a heating temperature of the steel slab (slab) of lower than 1,000°C results in insufficient dissolution of carbides to reduce the amount of steel that is subjected to solid solution strengthening by solute C and so forth, resulting in a failure to achieve the required strength.
  • a heating temperature of the steel slab of higher than 1,250°C results in excessively coarse crystal grains to deteriorate the fatigue crack growth resistance.
  • the heating temperature of the steel slab is 1,000°C to 1,250°C.
  • the heating temperature of the steel slab is preferably 1,030°C or higher.
  • the heating temperature of the steel slab is preferably 1,200°C or lower.
  • the steel slab (slab) is heated to the heating temperature all the way to the center.
  • the recrystallization of crystal grains is required to be promoted during hot rolling in the recrystallization temperature range to inhibit the formation of coarse grains.
  • the total rolling reduction in the recrystallization temperature range is less than 35%, recrystallization is insufficient, resulting in coarse grains remaining.
  • the total rolling reduction in the recrystallization temperature range is 35% or more, preferably 38% or more.
  • the total rolling reduction in the recrystallization temperature range is more than 55%, coarsening of the crystal grains can be inhibited.
  • the rolling reduction in the non-recrystallization range is insufficient, resulting in a failure to refine the crystal grains.
  • the total rolling reduction in the recrystallization temperature range is 55% or less, preferably 52% or less.
  • the lower limit temperature Tnr (°C) of the recrystallization temperature range can be calculated, for example, from the components of the steel using the following formula.
  • the temperature in the hot rolling is the surface temperature of the material being rolled (steel slab or steel plate), and the surface temperature can be measured with a radiation thermometer or the like.
  • Tnr C ⁇ 174 ⁇ log % Nb % C + 12 / 14 % N + 1444 where in the above formula, each [%X] is the element X content (mass%) of the steel.
  • Rolling in the partial recrystallization range needs to be started in a uniform grain state free of coarse grains by ensuring a sufficient rolling reduction in the final rolling pass in the recrystallization temperature range to sufficiently promote recrystallization, in addition to the total rolling reduction in the recrystallization temperature range of 35% or more and 55% or less.
  • the rolling reduction in the final rolling pass in the recrystallization temperature range is less than 10%, recrystallization is insufficient, and the grains grow into coarse grains during the holding time between rough rolling and the start of finish rolling.
  • the rolling reduction of the final rolling pass in the recrystallization temperature range is 10% or more, preferably 11% or more.
  • the upper limit of the rolling reduction in the final rolling pass in the recrystallization temperature range is not limited, and a higher rolling reduction is more preferred. As an example, the rolling reduction in the final rolling pass in the recrystallization temperature range is 20% or less.
  • the rolling reduction in the final rolling pass in the temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range is 15% or more, preferably 16% or more.
  • the upper limit of the rolling reduction in the final rolling pass in the above temperature range is not particularly limited, and a higher rolling reduction is more preferred. As an example, the rolling reduction in the final rolling pass in the above temperature range is 25% or less.
  • Rolling at a temperature lower than is effective in refining grains because more strain is introduced when the rolling is performed at a lower temperature. For this reason, rolling is preferably performed at a low temperature lower than (the lower limit temperature of the recrystallization temperature range - 80°C) within a range in which the cooling start temperature can be observed.
  • the cooling start temperature in the cooling step after the hot rolling step needs to be higher than or equal to the Ar 3 transformation point in terms of a surface temperature of the steel plate.
  • the finish rolling temperature needs to be set.
  • the term "Ar 3 transformation point" refers to the ferrite transformation start temperature during cooling, and can be calculated, for example, from the components of the steel using the following formula.
  • the surface temperature of the steel plate can be measured using a radiation thermometer or the like.
  • Cooling Start Temperature Temperature Higher than or Equal to Ar 3 Transformation Point (°C) in Terms of Surface Temperature of Steel Plate
  • the steel plate After the hot rolling step, the steel plate is cooled (controlled cooling).
  • the surface temperature of the steel plate at the start of cooling is lower than the Ar 3 transformation point (°C)
  • ferrite forms before cooling, resulting in a significant decrease in strength.
  • the surface temperature of the steel plate at the start of cooling is higher than or equal to the Ar 3 transformation point (°C).
  • the surface temperature of the steel plate at the start of cooling is the temperature of the surface region of the steel plate where the cooling start temperature is the lowest.
  • the surface temperature of the steel plate at the start of cooling is the surface temperature of the steel plate at its trailing end portion when, for example, the steel plate is cooled while traveling in one direction with respect to a cooling device.
  • the surface temperature of the steel plate is the surface temperature of the region that is cooled last.
  • the upper limit of the surface temperature of the steel plate at the start of cooling is the above-mentioned finish rolling temperature.
  • the difference in cooling start time for the entirety of the steel plate is more than 50 seconds, the difference in the temperature of the steel plate is large, resulting in a large variation in the temperature of the steel plate when cooling is stopped. This increases the variation in Vickers hardness at 0.25 mm below the surface of the steel plate and deteriorates the HISC resistance. Therefore, the difference in cooling start time for the entirety of the steel plate is within 50 seconds, preferably within 45 seconds. Specifically, for example, when the steel plate is cooled while traveling in one direction with respect to the cooling device, the difference in cooling start time between the leading end of the steel plate and the trailing end of the steel plate is within 50 seconds.
  • the difference in cooling start time between the first region and the last region is within 50 seconds.
  • the difference in cooling start time for the entire steel plate may be 0 seconds.
  • the temperature range of 750°C to 550°C is an important temperature range for bainite transformation. It is thus important to control the cooling rate in this temperature range.
  • the average cooling rate in the temperature range is more than 50 °C/s, the hardness may vary, and the HISC resistance after pipe production deteriorates. Therefore, the average cooling rate is 50 °C/s or less, preferably 45 °C/s or less. When the cooling rate is too low, ferrite and pearlite are formed, resulting in insufficient strength.
  • the average cooling rate in the above-mentioned temperature range is 15 °C/s or more, preferably 17 °C/s or more.
  • the average cooling rate in the temperature range of 550°C to the cooling stop temperature in terms of the steel plate temperature at 0.25 mm below the surface of the steel plate is preferably 150 °C/s or more.
  • the average cooling rate is preferably 250 °C/s or less from the viewpoint of more easily inhibiting the variation in hardness.
  • the average cooling rate from 750°C to 550°C at the center of the plate thickness is less than 15 °C/s, a granular bainite structure is not formed, resulting in a reduction in strength. Therefore, the average cooling rate from 750°C to 550°C at the center of the plate thickness is 15 °C/s or more.
  • the average cooling rate is preferably 17 °C/s or more from the viewpoint of inhibiting the variation of the microstructure.
  • the average cooling rate is 50 °C/s or less, preferably 45 °C/s or less, in order to inhibit a variation in grain size.
  • the cooling in the temperature range of 550°C or lower in terms of the steel plate temperature at the center of the plate thickness is not particularly limited, but the average cooling rate in the above-mentioned temperature range is preferably 15 °C / s or more from the viewpoint of inhibiting the variations of the microstructure and grain size. From the above viewpoint, the average cooling rate in the above temperature range is preferably 50 °C/s or less.
  • the steel plate temperature at 0.25 mm below the surface of the steel plate and at the center of the plate thickness cannot be measured directly and physically. However, based on the surface temperature at the start of cooling measured by a radiation thermometer and a target surface temperature at the stop of cooling, the temperature distribution in the thickness section is calculated by difference calculation using, for example, a process computer, and the temperature distribution can be obtained in real time from the results.
  • the temperature at 0.25 mm below the surface of the steel plate in the temperature distribution is defined as the "temperature of the steel plate at 0.25 mm below the surface of the steel plate” in the present specification.
  • the temperature at the center of the plate thickness in the temperature distribution is defined as the "steel plate temperature at the center of the plate thickness" in the present specification.
  • Cooling Stop Temperature 250°C to 550°C in Terms of Steel Plate Temperature at 0.25 mm below Surface of Steel Plate and at Center of Plate Thickness
  • the cooling stop temperature is 550°C or lower, preferably 500°C or lower.
  • the cooling stop temperature is 250°C or higher, preferably 300°C or higher.
  • a steel pipe for transporting hydrogen (UOE steel pipe, electric resistance welded steel pipe, spiral steel pile, or the like) suitable for transporting high-pressure hydrogen gas can be manufactured by forming the high strength steel plate of the present invention into a tubular shape using press bend forming, roll forming, UOE forming, or the like, and then welding the butting portions.
  • the high-pressure hydrogen refers to, for example, a hydrogen gas environment with a pressure of 15 MPa or more.
  • a UOE steel pipe is manufactured by groove cutting the ends of a steel plate, forming the steel plate into a steel pipe shape by C press, U press, and O press, then seam welding the butting portions by inner surface welding and outer surface welding, and optionally subjecting it to an expansion process.
  • Any welding method may be used as long as sufficient joint strength and joint toughness can be obtained, but it is preferable to use submerged arc welding from the viewpoint of excellent welding quality and production efficiency.
  • a steel pipe produced by forming a steel plate into a pipe shape using press bend forming and then seam welding the butting portions may also be subjected to pipe expansion.
  • Steels having the chemical compositions given in Table 1 were formed into steel slabs (slabs) by a continuous casting method.
  • the steel slabs were heated to the heating temperature given in Table 2 and then subjected to hot rolling and cooling under the conditions given in Table 2 to form steel plates having the final thickness given in Table 2.
  • each steel plate was subjected to controlled cooling using a water cooling-type controlled cooling device while the steel plate was traveling in one direction. Thereafter, the edges of the steel plate were subjected to groove cutting, and the steel plate was formed into a steel pipe shape by C press, U press, and O press.
  • the Vickers hardness (HV 0.5) was measured at 100 equally spaced points in the plate width direction from each of the leading end and the trailing end in the rolling direction of the steel plate in accordance with JIS Z 2244 (2009), the 100 equally spaced points being located at 0.25 mm below the surface of the steel plate.
  • the average value and the standard deviation ⁇ of the Vickers hardness values (HV 0.5) at 200 points in total were determined.
  • the leading end of the steel plate in the rolling direction is a position 1 m downstream of the leading edge of the steel plate in the rolling direction.
  • the trailing end of the steel plate in the rolling direction is a position 1 m upstream of the trailing edge of the steel plate in the rolling direction.
  • Table 3 presents the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3 ⁇ .
  • a sample for observing the metallic microstructure was taken from the central portion of the width of the steel plate obtained as described above. A section of this sample perpendicular to the plate width direction was mirror-polished and then etched with colloidal silica. Thereafter, crystal data was collected by an electron backscatter diffraction (EBSD) method in a field of view of 1 mm ⁇ 1 mm at the center of the plate thickness (measurement step: 0.8 ⁇ m). After the data collection, a boundary having an orientation difference of 15° or more was determined as a crystal grain boundary using OIM-Analysis (OIM Analysis software, manufactured by EDAX). The equivalent circular diameter was calculated as the crystal grain size from the area of each crystal grain. A frequency distribution table was prepared for all crystal grains measured, and the crystal grain size corresponding to 20% of the cumulative relative frequency from the largest crystal grain size calculated was defined as the "top 20% grain size". The measurement results are presented in Table 3.
  • a CT test specimen conforming to ASTM E 647 was taken from the steel plate obtained as described above in such a manner that the direction of load application was parallel to the rolling direction.
  • the CT test specimen was a 10-mm-thick test specimen taken from the 1/2 thickness position.
  • the length of the fatigue crack was measured by the compliance method using a clip gage, and the fatigue crack growth rate in a high-pressure hydrogen gas of 21 MPa was determined.
  • the fatigue crack growth rate (mm/cycle) when the stress intensity factor range ⁇ K was 45 (MPa ⁇ m 1/2 ) was evaluated. Table 3 presents the results.
  • a full-thickness test piece in the direction perpendicular to the rolling direction was used as a test piece for a tensile test, and the tensile test was performed in accordance with the provisions of JIS Z2241 (2011) to measure tensile strength and yield strength. Table 3 presents the results.
  • a test piece (coupon) cut out from the obtained steel pipe was flattened, and then a test specimen of 3 mm ⁇ 10 mm ⁇ 50 mm was taken from the inner surface of the steel pipe.
  • a test specimen containing both the weld and the base material was collected in addition to a test specimen containing only the base material without the weld.
  • the inner surface to be tested was left intact without removing the scale in order to leave the state of the outermost layer. That is, the position 0.25 mm below the surface of the steel plate is included in the test specimen.
  • the target ranges of the present invention are as described below.
  • the average value of the Vickers hardness at 0.25 mm below the surface of the plate + 3 ⁇ is 225 HV or less.
  • the top 20% grain size in the microstructure at the center of the plate thickness is 30 ⁇ m or less.
  • the fatigue crack growth rate is less than 2.0 ⁇ 10 -2 (mm/cycle) when the stress intensity factor range ⁇ K is 45 (MPa ⁇ m 1/2 ).
  • the tensile strength is 535 MPa or more.
  • no cracking is observed in the evaluation of the HISC resistance (four-point bending test). [Table 2] No.
  • No. 1 to No. 9 and No. 33 to No. 35 are Examples in which the chemical composition and production conditions satisfy the appropriate ranges of the present invention.
  • Table 3 in each of No. 1 to No. 9 and No. 33 to No. 35, the average value of the Vickers hardness + 3 ⁇ was 225 HV or less at 0.25 mm below the surface of the steel plate as a high strength steel plate.
  • the top 20% grain size in the microstructure at the center of the plate thickness was less than 30 ⁇ m.
  • the fatigue crack growth rate was less than 2.0 ⁇ 10 -2 (mm/cycle) when the stress intensity factor range ⁇ K was 45 (MPa ⁇ m 1/2 ).
  • the tensile strength was 535 MPa or more.
  • the HISC resistance was also good.
  • the chemical composition of the steel plate of each of No. 10 to No. 20 is outside the scope of the present invention.
  • the solid solution strengthening was insufficient, resulting in insufficient strength.
  • the Vickers hardness at 0.25 mm below the surface of the steel plate was increased, resulting in poor HISC resistance.
  • the grain growth was not sufficiently inhibited by precipitates, resulting in poor fatigue crack growth resistance.
  • No. 21 to No. 32 are Comparative Examples in which the chemical compositions are within the scope of the present invention, but the production conditions are outside the scope of the present invention.
  • No. 21 since the heating temperature of the steel slab (slab) was low, carbides were not sufficiently dissolved, resulting in low strength.
  • No. 22 since the heating temperature of the steel slab was high, the crystal grains were coarsened, resulting in a deterioration in fatigue crack growth resistance.
  • No. 23 since the total rolling reduction in the recrystallization temperature range was insufficient, coarse grains remained, resulting in a deterioration in fatigue crack growth resistance.
  • No. 21 to No. 32 are Comparative Examples in which the chemical compositions are within the scope of the present invention, but the production conditions are outside the scope of the present invention.
  • No. 21 since the heating temperature of the steel slab (slab) was low, carbides were not sufficiently dissolved, resulting in low strength.
  • No. 22 since the heating temperature of the steel slab was high, the crystal grains were coarsened
  • the high strength steel plate for a hydrogen transport steel pipe can be provided, the high strength steel plate having excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment.

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EP23839276.5A 2022-07-14 2023-05-11 Hochfestes stahlblech für ein wasserstofftransportstahlrohr, herstellungsverfahren dafür und wasserstofftransportstahlrohr Pending EP4520845A4 (de)

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JP6825748B2 (ja) * 2018-09-28 2021-02-03 Ｊｆｅスチール株式会社 耐サワーラインパイプ用高強度鋼板およびその製造方法並びに耐サワーラインパイプ用高強度鋼板を用いた高強度鋼管
US12351896B2 (en) * 2018-12-26 2025-07-08 Jfe Steel Corporation Steel material for high-pressure hydrogen gas environment, steel structure for high-pressure hydrogen gas environment, and methods for producing steel material for high-pressure hydrogen gas environment
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