EP3042976A1 - Stahlblech für hochfestes dickwandiges leitungsrohr mit aussergewöhnlichem aussäuerungswiderstand, quetschfestigkeitseigenschaften und duktilität bei niedrigen temperaturen sowie leitungsrohr - Google Patents

Stahlblech für hochfestes dickwandiges leitungsrohr mit aussergewöhnlichem aussäuerungswiderstand, quetschfestigkeitseigenschaften und duktilität bei niedrigen temperaturen sowie leitungsrohr Download PDF

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Publication number
EP3042976A1
EP3042976A1 EP14840842.0A EP14840842A EP3042976A1 EP 3042976 A1 EP3042976 A1 EP 3042976A1 EP 14840842 A EP14840842 A EP 14840842A EP 3042976 A1 EP3042976 A1 EP 3042976A1
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EP
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Prior art keywords
less
steel plate
low
strength
thick
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EP14840842.0A
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English (en)
French (fr)
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EP3042976A4 (de
EP3042976B1 (de
Inventor
Takuya Hara
Taishi Fujishiro
Yasuhiro Shinohara
Eiji Tsuru
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Nippon Steel Corp
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Nippon Steel and Sumitomo Metal Corp
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • C21D9/14—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes wear-resistant or pressure-resistant pipes
    • 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
    • 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
    • C21D1/78—Combined heat-treatments not provided for above
    • 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
    • C21D8/0247—Modifying the physical properties 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 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
    • 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
    • C22C38/18—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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • 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
    • C21D2221/00—Treating localised areas of an article
    • C21D2221/10—Differential treatment of inner with respect to outer regions, e.g. core and periphery, respectively
    • 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
    • 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
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C21D8/105—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies of ferrous alloys

Definitions

  • the present invention relates to steel plate for thick-gauge high-strength linepipe which is excellent in sour resistance, collapse resistance, and low-temperature toughness, in particular steel plate for thick-gauge high-strength linepipe which is excellent in sour resistance, collapse resistance, and low-temperature toughness which is therefore optimal for linepipe for transport of natural gas or crude oil and relates to linepipe having excellent sour resistance, collapse resistance, and low-temperature toughness.
  • PLT 1 proposes the method of rolling in a temperature region where the microstructure becomes the dual phase of austenite and ferrite (dual phase region). According to this method, it is possible to make the microstructure of a thick-gauge material into a fine acicular ferrite structure in which island martensite is mixed.
  • the invention which is disclosed in PLT 2 considers how to improve crack propagation resistance and low-temperature toughness, but does not consider how to improve sour resistance and collapse resistance. Further, the invention which is disclosed in PLT 3 considers low-temperature toughness and collapse resistance, but does not consider how to improve sour resistance. Further, the invention which is disclosed in PLT 4 attempts to achieve a balance of compressive strength and low-temperature toughness and both high compressive strength and sour resistance, but does not consider the above-mentioned collapse resistance (0.2% flow stress of compression in circumferential direction after aging at 200°C).
  • PLT 5 proposes a process of production which lowers the content of C, makes the microstructure into a low temperature transformation microstructure which is formed mainly of bainite, and, based on this steel material whose toughness is improved, adds Mo to improve the hardenability and keeps down the addition of Al so as to make use of the bainite in the grains.
  • the invention which is disclosed in PLT 5 improves the hardenability of the base material and makes the effective grain size of the HAZ finer by composing the steel plate as a whole of uniform microstructure formed of mainly bainite.
  • the invention which is disclosed in PLT 5 is aimed at improving the low-temperature toughness of the weld zone and does not consider how to improve the sour resistance and collapse resistance.
  • steel plate for linepipe often had a plate thickness of a thin 20 mm or less. If a strength of the X65 class or so of the API standard, it was possible to easily secure various characteristics such as the sour resistance, low-temperature toughness, and collapse resistance. This was because with hot-rolling, the reduction rate was sufficiently secured and the effective grain size became finer and, further, the difference in cooling rate between the surface layers and mid-thickness portion due to accelerated cooling was small so the microstructure became uniform. In this regard, if the plate thickness is 25 mm or more, in particular 30 mm or more, it becomes difficult to satisfy all of the requirements of sour resistance, low-temperature toughness, and collapse resistance.
  • the present invention in consideration of this situation, has as its object the provision of thick-gauge high-strength linepipe which is optimal as a material for linepipe for transport of natural gas or crude oil and has a good balance of sour resistance, collapse resistances, and low-temperature toughness and steel plate for the thick-gauge high-strength linepipe.
  • the inventors engaged in intensive studies focusing on the microstructure and crystal grain size in steel plate for linepipe so as to obtain steel plate for thick-gauge high-strength linepipe which is excellent in sour resistance, collapse resistance, and low-temperature toughness.
  • thick-gauge linepipe also referred to as "thick-gauge steel pipe”
  • the compositions, microstructures, processes of production, etc. for achieving (1) both strength and sour resistance, (2) both strength and collapse resistance of thick-gauge steel pipe, and (3) both strength and low-temperature toughness of thick-gauge steel pipe can be summarized as follows:
  • the inventors studied the method of making use of the feature of being thick-gauged, that is, using hot-rolling and subsequent accelerated cooling, to control the structure by utilizing the temperature difference between the surfaces and the center part due to the plate thickness. Further, they took note of the fact that at the center part of plate thickness, securing the sour resistance is extremely important while at the surface layers, securing the collapse resistance is extremely important. Further, to secure the low-temperature toughness, they studied refinement of the effective grain size.
  • the structure of the surface layers should be made a structure in which deformed ferrite with an area percentage of 5% or more should be formed so as to satisfy the collapse resistance and should be suppressed in MA and given a balance of one or both of polygonal ferrite and bainite so as to secure low-temperature toughness.
  • the present invention investigated the relationship between the allowed amount of deformed ferrite and plate thickness and discovered the optimum relationship.
  • the present invention was made based on these discoveries and has as its gist the following:
  • steel plate for thick-gauge high-strength linepipe which is excellent in sour resistance, collapse resistance, and low-temperature toughness which has a gauge thickness of 25 to 45 mm and, after formation into pipe, a YS of 440 MPa or more, TS of 500 to 700 MPa, DWTT shear area at -10°C of 85% or more, and compressive strength in the circumferential direction after aging at 200°C (0.2% flow stress) of 450 MPa or more.
  • the contribution to industry is extremely remarkable.
  • the steel plate for thick-gauge high-strength linepipe excellent in sour resistance, collapse resistance, and low-temperature toughness of the present invention (below, also simply referred to as "steel plate for linepipe” or “steel plate”) and a method of production of the same will be explained.
  • the reasons for limitation of the components in the steel plate for thick-gauge high-strength linepipe of the present embodiment base material of linepipe
  • the symbols % mean mass% unless otherwise indicated.
  • C is an element which improves the strength of steel plate.
  • 0.04% or more has to be added.
  • 0.05% or more more preferably 0.055% or more of C is added.
  • the upper limit of the amount of C is made 0.08%.
  • the upper limit of the amount of C is made 0.07%, more preferably the upper limit is made 0.065%.
  • Mn is an element which contributes to improvement of the strength and toughness of steel plate.
  • Mn is an element which contributes to improvement of the strength and toughness of steel plate.
  • 1.2% or more of Mn is added to secure the strength of the steel plate.
  • 1.4% or more, more preferably 1.5% or more of Mn is added.
  • the upper limit of the amount of Mn is made 2.0% or less.
  • the upper limit of the amount of Mn is made 1.8% or less, more preferably 1.7% or less.
  • Nb is an element which forms carbides and nitrides and contributes to the improvement of strength. Further, it suppresses recrystallization and promotes grain refinement during hot-rolling. For that reason, the lower limit of the amount of Nb is made 0.005% or more. Preferably, the lower limit of the amount of Nb is made 0.010% or more, more preferably 0.015% or more. On the other hand, if Nb is excessively added, the strength excessively rises and the low-temperature toughness is impaired, so the upper limit of the amount of Nb is made 0.05% or less. Preferably, the upper limit of the amount of Nb is made 0.04% or less, more preferably 0.03% or less.
  • Ti is an element which forms nitrides and exerts an effect on the grain refinement of the microstructure.
  • the lower limit of the amount of Ti is made 0.005% or more to make the effective grain size finer.
  • the lower limit of the amount of Ti is made 0.008% or more, more preferably 0.01% or more.
  • the upper limit of the amount of Ti is made 0.03% or less.
  • the upper limit of the amount of Ti is made 0.02% or less, more preferably 0.015%.
  • Ca is an element which controls the form of sulfides and improves the sour resistance.
  • the lower limit of the amount of Ca is made 0.0005% or more.
  • the lower limit of the amount of Ca is made 0.0010%, more preferably 0.0015%.
  • the upper limit of the amount of Ca is made 0.0050%.
  • the upper limit of the amount of Ca is made 0.0040% or less, more preferably 0.0030% or less.
  • N In the present embodiment, nitrides are utilized to make the microstructure of the steel finer, so the content of N is made 0.001% or more. Preferably, the amount of N is made 0.002% or more, more preferably 0.003% or more. On the other hand, if N is excessively contained, coarse nitrides will be formed and the low-temperature toughness will be impaired, so the upper limit of the amount of N is made 0.008%. Preferably, the upper limit of the amount of N is 0.007% or less, more preferably 0.006% or less.
  • Si and Al are deoxidizing elements. If added for the purpose of deoxidation, it is sufficient to use either one, but both may be used as well. Note that if Si and Al are excessively added, they impair the characteristics of the steel plate, so in the present embodiment, the upper limits of the contents of Si and Al are made the following:
  • Si If Si is excessively added, hard MA is formed in particular at the heat affected zone (HAZ) and the toughness of the seam weld zone of the steel pipe is made to fall, so the upper limit of the amount of Si is made 0.5% or less. Preferably, the amount of Si is made 0.3% or less, more preferably 0.25% or less. Note that, as explained above, Si is an element which is used for deoxidation and is an element which contributes to the rise in strength as well, so preferably the lower limit of the amount of Si is 0.05% or more, more preferably 0.10% or more.
  • Al is a useful deoxidizing element.
  • the lower limit of the amount of Al is 0.001% or more, more preferably 0.003% or more.
  • the upper limit of the amount of Al is made 0.05% or less.
  • the upper limit of the amount of Al is made 0.04% or less, more preferably 0.03% or less. Further, by restricting the amount of Al to 0.005% or less, the HAZ toughness can be improved.
  • P, S, and O are contained as unavoidable impurities. If excessively contained, the characteristics of the steel plate are impaired, so in the present embodiment, the upper limits of the contents of P, S, and O are set as follows:
  • P is an element which causes embrittlement of the steel. If over 0.03% is contained, the low-temperature toughness of the steel is impaired, so the upper limit is made 0.03% or less. Preferably, the upper limit of the amount of P is made 0.02% or less, more preferably 0.01% or less.
  • S is an element which forms MnS and other sulfides. If over 0.005% is contained, the low-temperature toughness and the sour resistance are made to fall, so the upper limit is made 0.005% or less. Preferably, the amount of S is made 0.003% or less, more preferably 0.002%.
  • O If O is contained in over 0.005%, coarse oxides are formed and the low-temperature toughness of the steel is made to fall, so the upper limit of the content is made 0.005% or less. Preferably, the upper limit of the amount of O is made 0.003% or less, more preferably 0.002% or less.
  • one or more of Cu, Ni, Cr, Mo, W, V, Zr, Ta, and B can be added.
  • Cu is an element which is effective for making the strength rise without making the low-temperature toughness fall.
  • 0.01% or more of Cu is added, more preferably 0.1% or more is added.
  • Cu is an element which makes cracking occur more easily at the time of heating the steel slab or at the time of seam welding the steel pipe, so the amount of Cu is preferably made 0.50% or less. More preferably, the amount of Cu is made 0.35% or less, still more preferably 0.2% or less.
  • Ni is an element which is effective for improving the low-temperature toughness and strength.
  • 0.01% or more of Ni is added, more preferably 0.1% or more is added.
  • Ni is an expensive element. From the viewpoint of economy, the amount of Ni is preferably made 0.50% or less. More preferably, the amount of Ni is made 0.35% or less, still more preferably 0.2% or less.
  • Cr is an element which improves the strength of the steel by precipitation strengthening.
  • 0.01% or more of Cr is added, more preferably 0.1% or more is added.
  • the upper limit of the amount of Cr is preferably made 0.50% or less. More preferably, the amount of Cr is made 0.35% or less, still more preferably 0.2% or less.
  • Mo is an element which improves the hardenability and which forms carbonitrides to improve the strength.
  • 0.01% or more of Mo is added, more preferably 0.05% or more is added.
  • the upper limit of the amount of Mo is preferably made 0.50% or less. More preferably, the amount of Mo is made 0.2% or less, more preferably 0.15% or less.
  • W W, like Mo, is an element which improves the hardenability and which forms carbonitrides to improve the strength.
  • 0.0001% or more of W is added, more preferably the amount of W is made 0.01% or more, still more preferably 0.05% or more is added.
  • the upper limit of the amount of W is preferably made 0.50% or less. More preferably, the amount of W is made 0.2% or less, more preferably 0.15% or less.
  • V is an element which forms carbides or nitrides and which contributes to the improvement of strength.
  • 0.001% or more of V is added, more preferably 0.005% or more is added.
  • the amount of V is preferably made 0.10% or less. More preferably, the amount of V is made 0.05% or less, more preferably 0.03% or less.
  • Zr and Ta are elements which form carbides or nitrides and contribute to the improvement of strength.
  • Zr and Ta are preferably added in 0.0001% or more, more preferably 0.0005% or more, still more preferably 0.001% or more is added.
  • the upper limits of the amount of Zr and the amount of Ta are preferably 0.050% or less. More preferably, the amounts are 0.030% or less.
  • B is an element which can cause an improvement in the hardenability by addition in a fine amount.
  • 0.0001% or more of B is preferably added.
  • 0.0003% or more of B is added.
  • the amount of B is preferably made 0.0020% or less. More preferably, the amount of B is made 0.0010% or less.
  • one or more of Mg, REM, Y, Hf, and Re may be added.
  • Mg is an element which contributes to improvement of the sour resistance or low-temperature toughness by control of the form of the sulfides or formation of fine oxides.
  • 0.0001% or more of Mg is added, more preferably 0.0005% or more, still more preferably 0.001% or more is added.
  • the amount of Mg is preferably made 0.010% or less. More preferably, the amount of Mg is made 0.005% or less, still more preferably 0.003% or less.
  • REM, Y, Hf, and Re form sulfides and suppress the formation of MnS elongated in the rolling direction, in particular, contribute to the improvement of the sour resistance.
  • REM, Y, Hf, and Re are all preferably added in 0.0001% or more, more preferably 0.0005% or more, still more preferably 0.0010% or more.
  • the upper limit is preferably made 0.0050% or less. More preferably, the amount is made 0.0030% or less.
  • the balance besides the above elements is substantially comprised of Fe.
  • Unavoidable impurities and other elements which do not harm the action or effect of the present invention may also be added in trace amounts.
  • Unavoidable impurities mean components which are contained in the raw materials or which enter in the process of production and refer to components which are not deliberately included in the steel.
  • Si, Al, P, S, O, N, Sb, Sn, Co, As, Pb, Bi, and H may be mentioned.
  • P, S, O, and N as explained above, have to be controlled to Si: 0.5% or less, Al: 0.05% or less, P: 0.03% or less, S: 0.005% or less, O: 0.005% or less, and N: 0.008% or less.
  • Sb, Sn, Co, and As can be contained in amounts of 0.1% or less, Pb and Bi in 0.005% or less, and H in 0.0005% or less as unavoidable impurities. However, if in the usual ranges, do not particularly have to be controlled.
  • the optionally added elements of Cu, Ni, Cr, Mo, W, V, Zr, Ta, B, Mg, REM, Y, Hf, and Re in the steel plate for thick-gauge high-strength linepipe according to the present invention can be contained as unavoidable impurities even if not deliberately included.
  • these elements do not have a detrimental effect on the present invention even if the amounts of the added elements are below the lower limit so long as the amounts of the added elements are below the upper limit of the content in the case of deliberate inclusion explained above, so do not pose problems.
  • the carbon equivalent Ceq of the following (formula 2) which is calculated from the contents of the C, Mn, Ni, Cu, Cr, Mo, and V (mass%), is preferably made 0.30 to 0.50.
  • the lower limit of Ceq is more preferably 0.32 or more, still more preferably 0.35 or more, to raise the strength.
  • the upper limit of the Ceq is more preferably 0.45 or less, still more preferably 0.43 or less, to raise the low-temperature toughness.
  • Ceq C+Mn/6+(Ni+Cu)/15+(Cr+Mo+V)/5 ... (formula 2)
  • the cracking susceptibility parameter Pcm of the following (formula 3), which is calculated from the contents of the C, Si, Mn, Cu, Cr, Ni, Mo, and V (mass%), is preferably 0.10 to 0.20.
  • the lower limit of Pcm raises the strength, so is more preferably 0.12 or more, still more preferably 0.14 or more.
  • the upper limit of the Pcm raises the low-temperature toughness, so is more preferably 0.19 or less, still more preferably 0.18 or less.
  • Pcm C+Si/30+(Mn+Cu+Cr)/20+Ni/60+Mo/15+V/10 ... (formula 3)
  • the steel plate of the present invention has a plate thickness of 25 mm or more, more preferably a 30 mm or more thickness, and is suitable as steel plate for thick-gauge linepipe (25 mm to 45 mm). Further, the steel plate of the present invention utilizes the temperature difference of hot-rolling or difference of cooling rate of the accelerated cooling at the surface layers and the mid-thickness portion to control the structure and differs in microstructure at the surface layers and the mid-thickness portion.
  • the surface layer portion of the steel plate is the portion of 0.9 mm to 1.1 mm from the surface of the steel plate in the thickness direction (that is, the region within 0.1 mm in the directions to both the front and back surfaces centered at the positions of 1 mm in the thickness directions from the surfaces of the steel plate), while the center part of the steel plate is the region within 1 mm in the directions to both the front and back surfaces from the center of plate thickness.
  • deformed ferrite is ferrite which is elongated by hot-rolling in the rolling direction. Compared with polygonal ferrite which is formed by cooling after rolling, the dislocation density is higher. This is effective for improvement of the collapse resistance.
  • An optical micrograph of the cross-section of a surface layer portion of the steel plate of the present invention is shown in FIG. 1 . Further, the dark gray parts are deformed ferrite. Such a part is shown by the arrow mark. The surface layer portion which is shown in FIG. 1 contains deformed ferrite in 9.3%.
  • the inventors discovered that it is possible to suppress the deformed ferrite at the center part to raise the low-temperature toughness.
  • the thickness of the steel plate becomes greater, the temperature difference between the surface layers and the center in wall thickness becomes larger.
  • the gauge thickness of the steel plate becomes greater, the amount of deformed ferrite which can be produced at the center part of plate thickness becomes smaller, while the amount of deformed ferrite which can be produced at the surface layer portion becomes greater. Therefore, the inventors investigated the relationship of the gauge thickness of the steel plate and the amount of deformed ferrite at the surface layer portion and discovered the optimal range.
  • FIG. 2 shows the relationship between plate thickness of steel plate with a plate thickness of 25 mm to 45 mm and the upper limit S fe1 of the area percentage of deformed ferrite at the surface layer portion.
  • the area percentage of the deformed ferrite at the surface layer portion of the steel plate has to be the following lower limit value or more and the upper limit value or less.
  • Lower limit value of area percentage of deformed ferrite at surface layer portion of steel plate 5%
  • the area percentage of the deformed ferrite exceeds the S fe1 %, the surface layers harden and the low-temperature toughness is impaired, so the area percentage of the deformed ferrite is made S fe1 % or less.
  • the area percentage of the deformed ferrite for obtaining a sour resistance, collapse resistance, and low-temperature toughness optimal for a material of linepipe for transporting natural gas or crude oil depends on the plate thickness.
  • the temperature difference in hot-rolling between the surface layers and the mid-thickness portion and the difference in cooling rates at accelerated cooling are easily affected by the plate thickness, so the area percentage of the deformed ferrite is considered to have dependency on the plate thickness.
  • the MA at the surface layer portion is restricted to an area percentage of 8% or less.
  • the area percentage of the MA at the surface layer portion is made 5% or less, more preferably 3% or less.
  • the balance besides the above deformed ferrite and MA is a microstructure composed of one or both of polygonal ferrite and bainite.
  • Polygonal ferrite is effective for improvement of the low-temperature toughness. It is easily formed at the surface layer portion and gradually decreases toward the mid-thickness portion. Bainite is effective for improvement of the strength.
  • the amount of it is minor at the surface layer portion and gradually increases toward the mid-thickness portion. This is because at the mid-thickness portion, compared with the surface layers, the rolling temperature in hot-rolling and the start temperature of accelerated cooling become higher.
  • the area percentage of deformed ferrite is restricted to 5% or less.
  • the area percentage of deformed ferrite is preferably made 3% or less, more preferably 0%.
  • the area percentage of MA is restricted to 5% or less.
  • the area percentage of MA at the mid-thickness portion is made 4% or less, more preferably is made 2% or less.
  • the balance besides the deformed ferrite and MA is a microstructure comprised of one or both of acicular ferrite and bainite.
  • Polygonal ferrite is effective for improving the low-temperature toughness, but impairs the sour resistance, so at the mid-thickness portion, the microstructure is preferably a uniform one comprised of one or both of acicular ferrite and bainite.
  • the microstructures of the above-mentioned surface layer portion and mid-thickness portion can be observed by an optical microscope.
  • the area percentages of the deformed ferrite and MA can be found by image analysis of the optical micrographs of the structures. Note that, at the MA, repeller etching is performed and the area percentage of the non-colored structures is found by image analysis.
  • the polygonal ferrite which is produced at the time of accelerated cooling is granular.
  • the deformed ferrite is elongated in the rolling direction. Further, the deformed ferrite is high in dislocation density, so is hardened more compared with the polygonal ferrite.
  • the deformed ferrite and polygonal ferrite can be differentiated by the ratio of the long axis and short axis (aspect ratio) or the hardness.
  • Acicular ferrite and bainite are lath structures and can be differentiated by the deformed ferrite and polygonal ferrite.
  • the size of the region surrounded by high angle grain boundaries of a difference of orientation of 15° or more is made smaller to improve the low-temperature toughness.
  • EBSD electron backscatter diffraction
  • the low-temperature toughness of steel plate is evaluated by measuring the effective grain size at the mid-thickness portion and finding the average value. Further, as the means for measuring the effective grain size of different microstructures, electron backscatter diffraction is employed.
  • the effective grain size is defined as the circle equivalent diameter found by analyzing the structure in the longitudinal direction of the steel plate after rolling by EBSD. Note that, at the surface layer portion, the size can be made smaller by utilizing deformed ferrite or polygonal ferrite, but at the mid-thickness portion, formation of deformed ferrite or polygonal ferrite ends up being suppressed, so the prior austenite grains can be made finer by hot-rolling.
  • the steel plate which is used for the linepipe is preferably made a plate thickness of 25 mm or more. Further, the steel plate preferably has a 500 MPa or more tensile strength.
  • the steel plate after pipe formation that is, the part of the steel pipe other than the weld zone and HAZ, for example, the part of the steel pipe from the seam part to 90° to 180° positions (positions at 3 o'clock to 6 o'clock from seam part) also similarly preferably has a 440 MPa or more yield stress and a 500 to 700 MPa or more tensile strength.
  • the plate thickness of the steel plate is more preferably 30 mm or more, still more preferably 35 mm or more.
  • low-temperature toughness of the linepipe When laying pipeline at arctic regions, low-temperature toughness of the linepipe is considered required.
  • the low-temperature toughness can be evaluated by the drop weight tear test (DWT test).
  • the DWTT shear area at -10°C of steel plate before pipe formation is preferably 85% or more.
  • the plate thickness of the steel plate is preferably made 45 mm or less and the tensile strength of the steel plate is preferably 700 MPa or less.
  • the strength of the steel plate after pipe formation tends to become higher than the strength of the steel plate before pipe formation, but the tensile strength of the steel pipe after formation is also preferably made 700 MPa or less.
  • the compressive strength in the circumferential direction after aging at 200°C (0.2% flow stress) is preferably 450 MPa or more.
  • the steel plate according to the present invention is given structures which differ at the surface layers and the mid-thickness portion by performing one or more passes of hot-rolling in the temperature region where the microstructure of the surface layers become dual phase of ferrite and austenite (dual phase region) and further performing the accelerated cooling after the hot-rolling by water cooling or other means under conditions whereby the temperature of the surfaces of the steel plate becomes 400°C or less and heat is recuperated after stopping thereof. If the steel plate is thick in gauge, the temperature of the surface layers at the time of hot-rolling falls from the temperature at the mid-thickness portion. At the mid-thickness portion, formation of ferrite is suppressed compared with the surface layers.
  • the stopping temperature of accelerated cooling becomes higher at the mid-thickness portion than at the surfaces. If setting a condition of accelerated cooling so that the temperature of the surfaces is recuperated after the accelerated cooling, the temperature of the center part of the steel plate after stopping the accelerated cooling can be made 400°C or more, hardening of the mid-thickness portion can be suppressed, and the sour resistance can be secured.
  • the average effective grain size of the surface layers and mid-thickness portion is made 20 ⁇ m or less.
  • the effective grain size becomes smaller.
  • the mid-thickness portion formation of deformed ferrite and polygonal ferrite ends up being suppressed, so the prior austenite grains have to be made smaller in size.
  • the process of production of the steel plate according to the present invention will be explained in order.
  • steel containing the above components is smelted in the steelmaking process, then is cast to obtain a steel slab.
  • the casting can be performed by an ordinary method, but from the viewpoint of productivity, continuous casting is preferable.
  • the obtained steel slab is heated, hot rolled, and cooled by accelerated cooling to produce steel plate.
  • the heating of the steel slab which is performed for hot-rolling is also referred to as "reheating” and the heating temperature of the steel slab at this time is also called the "reheating temperature”.
  • the reheating temperature of hot-rolling is made 1000°C or more so as to dissolve the carbides, nitrides, etc. which is formed in the steel slab in the steel. Further, by making the reheating temperature 1000°C or more, hot-rolling in the recrystallization region that is over 900°C (recrystallization rolling) is possible and the structure of the steel can be made finer. Note that, the upper limit of the reheating temperature is not prescribed, but to suppress coarsening of the effective grain size, the reheating temperature is preferably made 1250°C or less. Further, the reheating temperature is more preferably made 1200°C to secure the low-temperature toughness, more preferably 1150°C or less.
  • the hot-rolling according to the present embodiment is comprised of a rolling process in the recrystallization region that is over 900°C, rolling in the non-recrystallization region that is 900°C or less, and rolling in the temperature region where the temperature at the surface of the steel plate becomes a temperature resulting in a dual phase of austenite and ferrite (dual phase region) in that order.
  • the hot-rolling may be started right after extraction from the heating furnace performing the reheating, so the start temperature of the hot-rolling is not particularly prescribed.
  • the reduction ratio at the recrystallization region is the ratio of the plate thickness of the steel slab and the plate thickness at 900°C.
  • non-recrystallization region rolling hot-rolling is performed at the non-recrystallization region that is 900°C or less (non-recrystallization region rolling).
  • the reduction ratio at the non-recrystallization region rolling it is necessary to set the reduction ratio at the non-recrystallization region rolling to 3.0 or more and promote the transformation by accelerated cooling. More preferably, the reduction ratio at non-recrystallization rolling is set to 4.0 or more.
  • the reduction ratio of non-recrystallization rolling is the ratio of the plate thickness at 900°C divided by the plate thickness after the end of non-recrystallization rolling.
  • the rolling is performed in the temperature region (dual phase region) of the temperature of the surfaces of the steel plate wherein dual phase of austenite and ferrite are formed.
  • the surface temperature of the steel plate becomes the beginning temperature of ferrite transformation Ar 3 or less, but during the period of the start to the end of the dual phase rolling, the temperature of the mid-thickness portion of the steel plate is maintained so as to be higher than the temperature of the surfaces of the steel plate and over Ar 3 .
  • Such a temperature distribution can be realized by, for example, performing accelerated cooling for a short time and lowering the temperature at only the surface layers.
  • the number of passes is set to 1 or more and the reduction rate is set to from 0.1 to 40%.
  • the start temperature of the later performed accelerated cooling also becomes the dual phase region, so hardening of the mid-thickness portion can be suppressed and the low-temperature toughness can be improved.
  • the "reduction rate" is the amount of reduction of the steel plate due to rolling, that is, the value which is obtained by dividing the difference between the thickness of the steel plate before rolling and the thickness of the steel plate after rolling by the thickness of the steel plate before rolling and can be expressed by a percent (%) etc. Further, at the portions between the surface layers and the mid-thickness portion, formation of polygonal ferrite is promoted.
  • Ar 3 can be calculated from the contents of C, Si, Mn, Ni, Cr, Cu, ad Mo (mass%).
  • Ar 3 905-305C+33Si-92(Mn+Ni/2+Cr/2+Cu/2+Mo/2)
  • the C, Si, Mn, Ni, Cr, Cu, and Mo in the above formula show the contents (mass%) of the elements. Further, Ni, Cu, Cr, and Mo are elements which are selectively added in the present invention. When not deliberately added, the content is calculated as "0" in the formula.
  • the lower limit of the reduction rate in dual phase rolling is set to 0.1% or more so as to cause the formation of deformed ferrite elongated in the rolling direction.
  • the reduction rate of the dual phase rolling is set to 1% or more, more preferably 2% or more.
  • the upper limit of the reduction rate in dual phase rolling is set to 40% or less since it is difficult to secure a reduction rate at a low temperature where the deformation resistance becomes higher.
  • the reduction rate in dual phase rolling is made 30% or less, more preferably 20% or less, still more preferably less than 10%.
  • the end temperature of the dual phase rolling is set to 700°C or more as a temperature of the surfaces of the steel plate so that the deformed ferrite is not excessively formed. If the hot-rolling end temperature becomes less than 700°C, ferrite transformation occurs at the mid-thickness portion and, due to the deformed ferrite, the low-temperature toughness and sour resistance sometimes fall. Further, if the hot-rolling end temperature falls, sometimes the formation of ferrite causes C to concentrate at the austenite and the formation of MA to be promoted. On the other hand, when the hot-rolling end temperature is too high, if the accelerated cooling stop temperature is lowered, the mid-thickness portion sometimes hardens and the low-temperature toughness falls.
  • accelerated cooling is immediately started.
  • air-cooling is allowed while the steel is transported from the exit side of the rolling mill to the accelerated cooling apparatus.
  • the accelerated cooling stop temperature is set to a temperature within temperature range of 200 to 400°C at the surfaces of the steel plate. If stopping the accelerated cooling at a temperature where the surface of the steel plate exceeds 400°C, polygonal ferrite is formed at the mid-thickness portion and the sour resistance falls. On the other hand, if performing accelerated cooling until the temperature of the surfaces of the steel plate becomes less than 200°C, the mid-thickness portion hardens and the low-temperature toughness falls. After accelerated cooling, air-cooling is performed in that state.
  • the temperature of the surface layers of the steel plate recovers at the time of air cooling. Therefore, the temperature of the mid-thickness portion reaches 400°C or more, the hardness falls, and the low-temperature toughness and sour resistance can be improved.
  • the above process of production can be used to produce the steel plate for high-strength linepipe according to the present invention. Further, when using the steel plate for high-strength linepipe according to the present invention as a material, it is possible to produce steel pipe for thick-gauge high-strength linepipe which is excellent in sour resistance, collapse resistance, and low-temperature toughness. Note that, when producing steel pipe, it is preferable to employ the UOE process of shaping the steel plate for high-strength linepipe according to the present invention by C-pressing, U-pressing, and O-pressing. Alternatively, the JCOE process can be used to produce steel pipe using the steel plate for high-strength linepipe according to the present invention.
  • the thick-gauge high-strength linepipe according to the present invention is produced by forming the steel plate for high-strength linepipe according to the present invention into a pipe shape, then arc welding the abutting ends.
  • arc welding submerged arc welding is preferably employed from the viewpoints of the toughness of the weld metal and the productivity.
  • the collapse resistance of the thick-gauge, high-strength linepipe according to the present invention can be evaluated by taking compression test pieces in the circumferential direction from the steel pipes produced by the above-mentioned methods.
  • the "slab thickness” of Table 3-1 and Table 3-2 shows the thicknesses of the obtained steel slabs (mm).
  • the steel slabs were reheated and hot-rolled in the recrystallization region that is over 900°C.
  • the "heating temperature” of Table 3-1 and Table 3-2 shows that reheating temperature
  • the "transport thickness” of Table 3-1 and Table 3-2 shows the plate thickness at 900°C after hot-rolling in the recrystallization region and before the hot-rolling in the later explained non-recrystallization region that is 900°C or less.
  • the "reduction ratio in recrystallization region” of Table 3-1 and Table 3-2 shows the ratio of the slab thickness divided by the transport thickness.
  • the steel plate having the transport thickness was hot-rolled in the non-recrystallization region that is 900°C or less.
  • the "plate thickness” of Table 3-1 and Table 3-2 shows the plate thickness after hot-rolling in the non-recrystallization region and before the later explained dual phase rolling, while the "non-recrystallization reduction ratio" of Table 3-1 and Table 3-2 is the value obtained by dividing the value of the transport thickness by the plate thickness after the end of the non-recrystallization rolling.
  • the final hot-rolling process before accelerated cooling was performed.
  • the surface temperature of the steel plate at the time of end of the final hot-rolling process is shown by the “finishing end temperature (°C)" in Table 3-1 and Table 3-2.
  • the number of rolling operations performed at the time of the final hot-rolling process is shown by the "no. of ⁇ + ⁇ reduction passes” in Table 3-1 and Table 3-2, while the reduction rate of the steel plate by the final hot-rolling process is shown by the " ⁇ + ⁇ reduction rate (%)" in Table 3-1 and Table 3-2.
  • Test pieces were taken from the surface layer portion and mid-thickness portions of the steel plates of the obtained Nos. 1 to 46. These were examined for structure by an optical microscope to find the area percentage of deformed ferrite and the area percentage of MA and confirm the structure of the balance.
  • the area percentage of MA was measured using a test piece etched by repeller etching. Further, the average values of the effective grain sizes at the surface layers and mid-thickness portion were found by EBSD.
  • DWTT shear area of Steel Plate Further, a full-thickness DWT test piece having the length direction corresponding to the width direction of the steel plate was taken from the center part of plate width of the steel plate of each of the obtained Nos. 1 to 46. The DWT test was also performed based on the API standard 2000 at -10°C to measure the DWTT shear area.
  • the obtained Nos. 1 to 46 steel plates were used to form pipes by the UOE process and were welded at the inside and outside surfaces by the heat inputs shown in Table 5-1 and Table 5-2 by submerged arc welding so as to produce outside diameter 30 to 36 inch steel pipes (the steel plate numbers and steel pipe numbers correspond to each other).
  • test pieces were taken from the steel pipes and were measured for strength and subjected to compression tests.
  • the test pieces were processed from the 3 o'clock positions of the steel pipes, in which the seam weld zones was defined as 0 o'clock, so that the longitudinal directions of the tensile test pieces matched the longitudinal directions of the steel pipes.
  • the strengths of the steel pipes were measured based on ASTM E9-09 so as to measure the yield strengths and tensile strengths in the longitudinal directions of the linepipes.
  • the 0.5% underload yield strength was defined as the yield strength.
  • the compression test pieces which were used for the compression test of steel pipe were obtained by taking parts which has 22 mm diameter and 66 mm length below 3 mm from the inside surfaces of the steel pipes at the 6 o'clock positions of the steel pipes when defining the seam weld zone of the steel pipes as 0 o'clock.
  • the compression test was conducted based on ASTM E9-09. The compressive strength after aging at 200°C for 10 minutes (0.2% flow stress) was found.
  • HIC test samples of 20 mm width and 100 mm length were taken from the 3 o'clock and 6 o'clock positions of the steel pipe.
  • the HIC test pieces were taken so that the center parts of gauge thickness of the steel pipes became the test positions.
  • the HIC test was based on TM0284 of the NACE (National Association of Corrosion and Engineer) and was performed using as the test solution the Solution B.
  • the crack length ratio (CLR) was used for evaluation.
  • the characteristics of the steel plates are shown in Table 4-1 and Table 4-2, while the characteristics of the steel pipes are shown in Tables 5-1 and 5-2.
  • the steel plates of Nos. 1 to 28 show examples of the present invention.
  • the steel pipes which were produced using these steel plates have yield stresses of 440 MPa or more and tensile strengths of 500 to 700 MPa in range.
  • the steel plates had tensile strengths of 500 MPa or more and had DWTT shear areas at -10°C of 85% or more.
  • the steel pipes produced by forming these steel plates into pipe shapes and then butt welding them were good ones with CLR of 10% or less after HIC tests and results of compression tests of 450 MPa or more after strain aging at 200°C.
  • Steel Nos. 29 to 46 are comparative examples. Steel Nos. 29 to 40 have contents of chemical components outside the range of the present invention, while Steel Nos. 41 to 46 have microstructures outside the range of the present invention and have at least one of the strength, low-temperature toughness, collapse resistance, and sour resistance of an inferior level. Steel No. 29 has a small amount of C and falls in strength and collapse resistance. On the other hand, Steel No. 30 has a large amount of C, Steel No. 31 has a large amount of Si, and Steel No. 32 has a large amount of Mn. In each comparative example, the tensile strength excessively rises and the low-temperature toughness falls. Further, the Ar 3 of Steel No.
  • 35 to 39 are examples which have large contents of elements which contribute to the formation of carbides, nitrides, oxides, and sulfides and which fall in low-temperature toughness due to precipitates and inclusions.
  • Steel Nos. 41 and 42 are examples which respectively are insufficient in reduction rate in the recrystallization region and reduction rate in the non-recrystallization region, become large in effective grain size, and fall in low-temperature toughness.
  • Steel No. 43 has an end temperature of hot-rolling of 700°C or more, but is low in Ar 3 and is not rolled in the dual phase region in the present invention, so deformed ferrite is not formed at the surface layer, the mid-thickness portion hardens, and the low-temperature toughness falls.
  • accelerated cooling stop temperature is high, deformed ferrite and MA are excessively formed at the mid-thickness portion, and the strength falls. Further, the accelerated cooling is stopped at the temperature where the temperature of the surface of the steel plate exceeds 400°C, so polygonal ferrite is formed at the mid-thickness portion and the sour resistance falls.
  • Steel Nos. 45 and 46 are examples where the rolling end temperatures are low, deformed ferrite and MA are excessively formed at the surface layer portion and mid-thickness portions, and the low-temperature toughnesses and sour resistances fall. Table 1-1 Steel Plate No.
EP14840842.0A 2013-08-30 2014-08-29 Stahlblech für dickwandiges hochfestes leitungsrohr mit aussergewöhnlicher korrosionsbeständigkeit, quetschfestigkeitseigenschaften und duktilität bei niedrigen temperaturen sowie leitungsrohr Active EP3042976B1 (de)

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CN110088333A (zh) * 2016-12-22 2019-08-02 株式会社Posco 具有优异的表面部分nrl-dwt特性的超厚钢材及其制造方法
EP3561112A4 (de) * 2016-12-22 2019-10-30 Posco Ultradickes stahlmaterial mit hervorragenden nrl-dwt-eigenschaften des oberflächenteils und verfahren zur herstellung davon
CN110088333B (zh) * 2016-12-22 2021-09-17 株式会社Posco 具有优异的表面部分nrl-dwt特性的超厚钢材及其制造方法
US11634784B2 (en) 2016-12-22 2023-04-25 Posco Co., Ltd Ultra-thick steel material having excellent surface part NRL-DWT properties and method for manufacturing same
US11453933B2 (en) 2016-12-23 2022-09-27 Posco High-strength steel material having enhanced resistance to crack initiation and propagation at low temperature and method for manufacturing the same
US11572600B2 (en) 2017-12-24 2023-02-07 Posco Co., Ltd Structural steel having excellent brittle crack propagation resistance, and manufacturing method therefor
US11578379B2 (en) 2017-12-26 2023-02-14 Posco Cold-rolled steel sheet having excellent high-temperature properties and room-temperature workability
US11591677B2 (en) 2017-12-26 2023-02-28 Posco Co., Ltd High-strength structural steel material having excellent fatigue crack propagation inhibitory characteristics and manufacturing method therefor
JP2019174452A (ja) * 2018-03-27 2019-10-10 Ｊｆｅスチール株式会社 鋼管の耐圧潰特性の評価方法
EP4092149A4 (de) * 2020-01-17 2023-05-03 Nippon Steel Corporation Stahlblech und stahlrohr

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CN105143487B (zh) 2017-03-08
JP5776860B1 (ja) 2015-09-09
RU2637202C2 (ru) 2017-11-30
EP3042976A4 (de) 2017-05-10
EP3042976B1 (de) 2020-05-13
KR20150139950A (ko) 2015-12-14
RU2016106920A (ru) 2017-10-05
JPWO2015030210A1 (ja) 2017-03-02
CN105143487A (zh) 2015-12-09
KR101730756B1 (ko) 2017-04-26

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