WO2015151468A1 - 耐歪時効特性及び耐hic特性に優れた高変形能ラインパイプ用鋼材およびその製造方法ならびに溶接鋼管 - Google Patents
耐歪時効特性及び耐hic特性に優れた高変形能ラインパイプ用鋼材およびその製造方法ならびに溶接鋼管 Download PDFInfo
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- strain aging
- ferrite
- bainite
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 97
- 239000010959 steel Substances 0.000 title claims abstract description 97
- 230000032683 aging Effects 0.000 title claims abstract description 48
- 239000000463 material Substances 0.000 title claims abstract description 30
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 29
- 238000000034 method Methods 0.000 title abstract description 32
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 49
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 46
- 229910001563 bainite Inorganic materials 0.000 claims abstract description 41
- 238000011282 treatment Methods 0.000 claims abstract description 33
- 239000000203 mixture Substances 0.000 claims abstract description 19
- 239000002184 metal Substances 0.000 claims abstract description 8
- 229910052751 metal Inorganic materials 0.000 claims abstract description 8
- 238000001816 cooling Methods 0.000 claims description 59
- 238000003303 reheating Methods 0.000 claims description 22
- 238000005096 rolling process Methods 0.000 claims description 16
- 239000012535 impurity Substances 0.000 claims description 6
- 230000000630 rising effect Effects 0.000 claims description 4
- 229910052719 titanium Inorganic materials 0.000 claims description 4
- 229910052720 vanadium Inorganic materials 0.000 claims description 4
- 229910052804 chromium Inorganic materials 0.000 claims description 3
- 229910052750 molybdenum Inorganic materials 0.000 claims description 3
- 229910052758 niobium Inorganic materials 0.000 claims description 3
- 229910052799 carbon Inorganic materials 0.000 claims description 2
- 229910052802 copper Inorganic materials 0.000 claims description 2
- 229910052748 manganese Inorganic materials 0.000 claims description 2
- 229910052757 nitrogen Inorganic materials 0.000 claims description 2
- 229910052698 phosphorus Inorganic materials 0.000 claims description 2
- 238000000576 coating method Methods 0.000 abstract description 18
- RWSOTUBLDIXVET-UHFFFAOYSA-N Dihydrogen sulfide Chemical compound S RWSOTUBLDIXVET-UHFFFAOYSA-N 0.000 abstract description 12
- 229910000037 hydrogen sulfide Inorganic materials 0.000 abstract description 12
- 238000010438 heat treatment Methods 0.000 description 19
- 229910001566 austenite Inorganic materials 0.000 description 13
- 238000003466 welding Methods 0.000 description 12
- 230000000694 effects Effects 0.000 description 11
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- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 7
- 230000007423 decrease Effects 0.000 description 7
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- VNWKTOKETHGBQD-UHFFFAOYSA-N methane Chemical compound C VNWKTOKETHGBQD-UHFFFAOYSA-N 0.000 description 4
- 229910001562 pearlite Inorganic materials 0.000 description 4
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- FAPWRFPIFSIZLT-UHFFFAOYSA-M Sodium chloride Chemical compound [Na+].[Cl-] FAPWRFPIFSIZLT-UHFFFAOYSA-M 0.000 description 2
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- 150000001247 metal acetylides Chemical class 0.000 description 2
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- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 1
- LFQSCWFLJHTTHZ-UHFFFAOYSA-N Ethanol Chemical compound CCO LFQSCWFLJHTTHZ-UHFFFAOYSA-N 0.000 description 1
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 1
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
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Classifications
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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
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B23—MACHINE TOOLS; METAL-WORKING NOT OTHERWISE PROVIDED FOR
- B23K—SOLDERING OR UNSOLDERING; WELDING; CLADDING OR PLATING BY SOLDERING OR WELDING; CUTTING BY APPLYING HEAT LOCALLY, e.g. FLAME CUTTING; WORKING BY LASER BEAM
- B23K31/00—Processes relevant to this subclass, specially adapted for particular articles or purposes, but not covered by only one of the preceding main groups
- B23K31/02—Processes relevant to this subclass, specially adapted for particular articles or purposes, but not covered by only one of the preceding main groups relating to soldering or welding
- B23K31/027—Making tubes with soldering or welding
-
- 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
- C21D1/62—Quenching devices
- C21D1/673—Quenching devices for die 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
- 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
-
- 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
- C21D8/02—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
-
- 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
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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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
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
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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
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
- C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B23—MACHINE TOOLS; METAL-WORKING NOT OTHERWISE PROVIDED FOR
- B23K—SOLDERING OR UNSOLDERING; WELDING; CLADDING OR PLATING BY SOLDERING OR WELDING; CUTTING BY APPLYING HEAT LOCALLY, e.g. FLAME CUTTING; WORKING BY LASER BEAM
- B23K2101/00—Articles made by soldering, welding or cutting
- B23K2101/04—Tubular or hollow articles
- B23K2101/10—Pipe-lines
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B23—MACHINE TOOLS; METAL-WORKING NOT OTHERWISE PROVIDED FOR
- B23K—SOLDERING OR UNSOLDERING; WELDING; CLADDING OR PLATING BY SOLDERING OR WELDING; CUTTING BY APPLYING HEAT LOCALLY, e.g. FLAME CUTTING; WORKING BY LASER BEAM
- B23K2103/00—Materials to be soldered, welded or cut
- B23K2103/02—Iron or ferrous alloys
- B23K2103/04—Steel or steel alloys
-
- 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
- C21D2211/001—Austenite
-
- 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
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
Definitions
- the present invention relates to a steel product for a line pipe having a small material deterioration after coating treatment at 300 ° C. or less, a method for producing the same, and a welded steel pipe, and has excellent HIC resistance in a wet hydrogen sulfide environment at pH 5 or higher, and APIs X60 to X70. Grade steel for line pipes.
- line pipes used for transportation of natural gas and crude oil are required to have high strength in order to improve transportation efficiency by high-pressure operation. Even if large deformation occurs due to ice gouging or ground deformation, high deformability capable of preventing the occurrence of cracks is required. For example, in pipelines laid on a cold seabed or an earthquake zone where ice gouging occurs, in addition to high uniform elongation, a line pipe having a low yield ratio of 90% or less is required.
- the steel sheet is formed into a cold tube and the butt portion is welded, and then the outer surface of the steel pipe is usually coated from the viewpoint of corrosion protection.
- a strain age hardening phenomenon occurs due to processing strain at the time of pipe making and heating at the time of coating treatment, yield stress increases, and the yield ratio in the steel pipe becomes larger than the yield ratio in the steel plate.
- HIC resistance hydrogen-induced crack resistance
- Patent Document 1 discloses a heat treatment method in which quenching from a two-phase region of ferrite and austenite is performed between quenching and tempering. Yes.
- Patent Document 2 discloses that rolling of a steel material is completed at an Ar 3 temperature or higher, and the subsequent accelerated cooling rate and A method for achieving a low yield ratio by controlling the cooling stop temperature to obtain a two-phase structure of acicular ferrite and martensite is disclosed.
- Patent Documents 3 and 4 include fine precipitates of composite carbide containing Ti and Mo, or composite carbide containing any two or more of Ti, Nb, and V.
- a low-yield-ratio, high-strength, high-toughness steel pipe excellent in strain aging characteristics and a method for producing the same are disclosed.
- Patent Document 5 describes a method for achieving a low yield ratio, high strength, and high uniform elongation excellent in strain aging resistance of API 5L X70 or less without greatly increasing the amount of alloying elements added to steel.
- a method is disclosed in which reheating is performed immediately after cooling to obtain a three-phase structure of bainite, polygonal ferrite, and island martensite (MA).
- Patent Document 6 discloses a method of reducing the hardness difference between ferrite and bainite as a method for obtaining the HIC resistance of a steel material having a two-phase structure of ferrite and bainite of X65 or higher.
- JP-A-55-97425 Japanese Unexamined Patent Publication No. 1-176027 JP 2005-60839 A Japanese Patent Laid-Open No. 2005-60840 JP 2011-74443 A JP 2003-301236 A
- Patent Document 6 has excellent HIC resistance in a wet hydrogen sulfide environment with a pH of 3.3 or higher.
- material design such as cleaning of steel components adapted to a severe sour environment is excessive, and an increase in manufacturing cost is a problem.
- the present invention shows excellent HIC resistance in a wet hydrogen sulfide environment of pH 5 or higher, and has a low yield ratio even after coating treatment.
- API 5L X60 to X70 grade steel material for high deformability line pipe and production thereof The object is to provide a method and a welded steel pipe.
- the present inventors diligently studied a manufacturing process of an appropriate component composition and a steel material manufacturing method, particularly controlled rolling and accelerated cooling after controlled rolling, and obtained the following knowledge.
- A It is possible to improve the HIC resistance by adding an appropriate amount of Cu and not containing Mo or even if it is contained to 0.01% or less.
- B In the accelerated cooling process, the cooling start temperature is appropriately controlled, the cooling is stopped in the middle of the bainite transformation, that is, in the temperature region where untransformed austenite exists, and then the bainite transformation finish temperature (hereinafter referred to as Bf point).
- the metallographic structure of the steel sheet is uniformly formed as a martensite-martensite (hereinafter referred to as MA) which is a hard phase in the mixed phase of ferrite and bainite.
- MA martensite-martensite
- before and after strain aging treatment Since the solid solution C can be reduced by setting the cooling start temperature and the cooling stop temperature in the accelerated cooling to appropriate temperatures, an increase in the yield ratio after strain aging can be suppressed.
- the present invention has been made by further studying the above findings and is as follows.
- component composition by mass, C: 0.030 to 0.100%, Si: 0.01 to 0.50%, Mn: 0.5 to 2.5%, P: 0.015%
- N 0.007% or less
- the metal structure has ferrite, bainite, and island martensite,
- the area fraction of the island martensite is 0.5 to 5.0%
- the hardness difference between the ferrite and the bainite is 60 or more in terms of Vickers hardness
- the metal structure has ferrite, bainite, and island martensite, and the island martensite has an area fraction of 0.5 to 5%, and the ferrite and bainite Strain resistance with a uniform elongation of 9% or more and a yield ratio of 90% or less before and after the strain aging treatment at a temperature of 300 ° C. or less.
- Aging characteristics and HIC resistance Excellent high deformability line method for producing a steel pipe sexual.
- the strain aging characteristic in the present invention refers to a characteristic that can suppress an excessive increase in the yield ratio even when heat treatment is performed at a temperature of 300 ° C. or lower.
- the HIC resistance in the present invention refers to a characteristic in which hydrogen-induced cracking does not occur in a wet hydrogen sulfide environment having a pH of 5 or higher.
- high deformability means the characteristic which satisfy
- C 0.030 to 0.100%
- C is an element that contributes to precipitation strengthening as a carbide. If C is less than 0.030%, the formation of MA (island martensite) is insufficient, so that sufficient strength cannot be secured, and a predetermined amount of hardness difference between ferrite and bainite cannot be secured, resulting in an increased yield ratio. .
- the C content is specified to be 0.030 to 0.100%.
- the C content is 0.05% or more.
- the C content is 0.09% or less.
- Si 0.01 to 0.50% Si is contained for deoxidation. When Si is less than 0.01%, the deoxidation effect is not sufficient. When Si exceeds 0.50%, deterioration of toughness and weldability is caused. For this reason, the Si content is specified to be 0.01 to 0.50%. Preferably, the Si content is 0.01 to 0.3%.
- Mn 0.5 to 2.5% Mn is contained for strength and toughness. If Mn is less than 0.5%, the effect is not sufficient, and since MA (island martensite) is insufficient, the yield ratio increases. Therefore, the Mn content is 0.5% or more, and preferably 1.2% or more, more preferably 1.5% or more, from the viewpoint of lowering the yield ratio due to MA generation. On the other hand, if Mn exceeds 2.5%, toughness and weldability deteriorate. For this reason, Mn content is prescribed
- P 0.015% or less
- P is an unavoidable impurity element that deteriorates weldability and HIC resistance.
- the P content is specified to be 0.015% or less.
- the P content is 0.010% or less.
- S 0.002% or less S generally becomes MnS inclusions in steel and deteriorates the HIC resistance. For this reason, the smaller the better. Since there is no problem if S is 0.002% or less, the upper limit of the S content is specified to be 0.002%. Preferably, the S content is 0.0015% or less.
- Cu 0.20 to 1.00%
- Cu is an important element in the present invention, and suppresses the intrusion of hydrogen into the steel and contributes to improving the HIC resistance. However, if Cu is less than 0.20%, the effect is not sufficient, and if it exceeds 1.00%, weldability deteriorates. Therefore, the Cu content is specified to be 0.20 to 1.00%. Preferably, the Cu content is 0.25% or more. Preferably, the Cu content is 0.5% or less.
- Mo 0.01% or less (including 0) Mo causes an increase in yield ratio due to strain aging and deterioration of the HIC resistance. For this reason, Mo is not contained, or even if it is contained, it is regulated to 0.01% or less. Preferably, the Mo content is 0.005% or less.
- Nb 0.005 to 0.05%
- Nb improves toughness by refining the structure, further forms carbides, and contributes to an increase in strength.
- the Nb content is specified to be 0.005 to 0.05%.
- the Nb content is 0.01 to 0.05%.
- Ti 0.005 to 0.040% Due to the pinning effect of TiN, Ti suppresses austenite coarsening during slab heating, improves base metal toughness, further reduces solid solution N, and suppresses an increase in yield ratio due to strain aging. However, if Ti is less than 0.005%, the effect is not sufficient, and if it exceeds 0.040%, the toughness of the weld heat affected zone deteriorates. For this reason, the Ti content is specified to be 0.005 to 0.040%. Preferably, the Ti content is 0.005 to 0.02%.
- Al 0.10% or less Al is contained as a deoxidizer. If Al exceeds 0.10%, the cleanliness of the steel decreases and the toughness deteriorates. For this reason, Al content is prescribed
- N 0.007% or less
- N is an unavoidable impurity element that causes an increase in yield ratio due to strain aging and deterioration of the toughness of the weld heat affected zone.
- the upper limit of N content is prescribed
- the N content is 0.006% or less.
- Ni 0.02 to 0.50%
- Ni is an element that contributes to improving HIC resistance and is effective in improving toughness and increasing strength. If Ni is less than 0.02%, the effect is not sufficient, and even if it exceeds 0.50%, the effect is saturated, which is disadvantageous in terms of cost. For this reason, when Ni is contained, the Ni content is specified to be 0.02 to 0.50%. Preferably, the Ni content is 0.2% or more. Preferably, the Ni content is 0.4% or less.
- Cr 1.00% or less Cr is an effective element for obtaining sufficient strength even at low C.
- Cr exceeds 1.00%, weldability deteriorates.
- the upper limit of Cr content is prescribed
- it is 0.1 to 0.5% or less.
- V 0.10% or less V improves toughness by refining the structure, further forms carbides, and contributes to improvement in strength.
- V exceeds 0.10%, the toughness of the heat affected zone is deteriorated.
- V content is prescribed
- the V content is 0.005% or more.
- the V content is 0.05% or less.
- Ca 0.0050% or less
- Ca is an element effective for improving toughness by controlling the form of sulfide inclusions.
- the Ca content is prescribed
- the Ca content is 0.001% or more.
- the Ca content is 0.004% or less.
- B 0.0050% or less B is an element effective in increasing the strength and improving the toughness of the weld heat affected zone. If B exceeds 0.0050%, the weldability deteriorates. For this reason, when it contains, B content is prescribed
- the balance other than the above components in the steel material of the present invention is Fe and inevitable impurities.
- the content of elements other than the above is not a problem as long as the effects of the present invention are not impaired.
- the metal structure of the steel sheet of the present invention is mainly composed of a three-phase structure composed of ferrite, bainite, and island martensite.
- the phrase “consisting mainly of a three-phase structure composed of ferrite, bainite, and island martensite” refers to a multiphase structure in which the area fraction of ferrite, bainite, and island martensite is 90% or more in total. The remainder is a structure having a total area fraction of 10% or less selected from martensite (excluding island martensite), pearlite, retained austenite, and the like.
- the area fraction of island martensite is 0.5 to 5.0%. Thereby, the yield ratio of 90% or less before and after the strain age hardening treatment can be satisfied. If the area fraction of island martensite is less than 0.5%, it may be insufficient to achieve a low yield ratio. Further, if the area fraction of island martensite exceeds 5.0%, the base material toughness and the HIC resistance may be deteriorated. Since the island-like martensite is stable without being decomposed even at the heating temperature (up to 300 ° C.) during the coating process, it is possible to achieve a low yield ratio even after the coating process in the present invention.
- the area fraction of ferrite and bainite is not particularly limited, but from the viewpoint of low yield ratio and HIC resistance, the area fraction of ferrite is 10% or more, and the area fraction of bainite is 10% or more. Preferably there is.
- the hardness difference between ferrite and bainite is 60 or more in terms of Vickers hardness (HV).
- HV Vickers hardness
- the hardness difference is preferably 180 or less at HV. More preferably, the hardness difference is 150 or less in HV.
- Each metallographic structure can be obtained, for example, by observation with an optical microscope or a scanning electron microscope, and image processing of the obtained microstructural photograph of at least three fields of view to determine the type of structure and the area fraction of each phase. it can.
- MA island martensite
- etching with a 3% nital solution (nitral alcohol solution) followed by electrolytic etching and observation.
- the MA can determine the area fraction by observing the microstructure photograph obtained by observation with a scanning electron microscope and at least three visual fields.
- the hardness is a value measured by a Vickers hardness tester, and an arbitrary load can be selected in order to obtain an indentation of an optimum size inside each phase. It is desirable to measure the hardness of ferrite and bainite under the same load. In addition, considering the local components of the microstructure or variations due to measurement errors, the hardness is measured at at least 15 different positions for each phase, and the average hardness of each phase is determined as the hardness of ferrite and bainite. It is preferable to use it. As the hardness difference when using the average hardness, the absolute value of the difference between the average value of the hardness of ferrite and the average value of the hardness of bainite is used.
- the uniform elongation is set to 9% or more and the yield ratio is set to 90% or less before and after the strain aging treatment at a temperature of 300 ° C. or less. It is preferable that the uniform elongation is 10% or more and the yield ratio is 88% or less from the viewpoint of high deformability before and after the strain aging treatment at a temperature of 300 ° C. or less.
- Ar 3 points are calculated from the following equation.
- Ar 3 (° C.) 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo
- the element symbol indicates the content (% by mass) of each element, and 0 when not contained.
- Heating temperature 1000-1300 ° C If the heating temperature is less than 1000 ° C, the solid solution of the carbide is insufficient and the required strength cannot be obtained, and if it exceeds 1300 ° C, the base material toughness deteriorates. For this reason, the heating temperature is regulated to 1000 to 1300 ° C.
- Rolling end temperature Ar 3 points or more
- the cumulative rolling reduction in a temperature range of 900 ° C. or lower is 50% or more.
- Cooling start temperature for accelerated cooling (Ar 3 -50) to (Ar 3 +30) ° C.
- the cooling start temperature is (Ar 3 -50) ° C. or higher, preferably (Ar 3 -30) ° C. or higher.
- the cooling start temperature exceeds (Ar 3 +30) ° C.
- the area fraction of ferrite decreases and the hardness difference between ferrite and bainite decreases, which is insufficient to achieve a low yield ratio. Therefore, the cooling start temperature is (Ar 3 +30) ° C. or lower, and preferably (Ar 3 +25) ° C. or lower.
- Accelerated cooling rate 5 ° C / s or more If the cooling rate is less than 5 ° C / s, pearlite is generated during cooling, and sufficient strength and low yield ratio cannot be obtained. To do.
- the cooling rate is 8 ° C./s or more, more preferably 10 ° C./s or more.
- the cooling rate is 100 ° C./s or less, more preferably 60 ° C./s or less.
- Cooling stop temperature 450-650 ° C
- the cooling stop temperature for accelerated cooling is an important manufacturing condition.
- C-enriched untransformed austenite present after reheating is transformed into MA (island martensite) during subsequent air cooling. That is, it is necessary to stop the cooling in a temperature range where untransformed austenite during the bainite transformation exists. If the cooling stop temperature is less than 450 ° C., the bainite transformation is completed, so MA (island martensite) is not generated during air cooling, and a low yield ratio cannot be achieved.
- the cooling stop temperature for accelerated cooling is set to 450 to 650 ° C.
- the cooling stop temperature is 515 ° C or higher, more preferably 530 ° C or higher.
- the cooling stop temperature is 635 ° C. or lower, more preferably 620 ° C. or lower.
- reheating it is preferable to perform reheating immediately after accelerated cooling. This is because it is preferable to perform reheating from a state in which untransformed austenite exists. “Immediately” is preferably within 120 seconds after cooling is stopped, from the viewpoint of reducing manufacturing efficiency and fuel cost required for heat treatment.
- This process of reheating to a temperature of 550 to 750 ° C. immediately after stopping the accelerated cooling at a temperature rising rate of 0.5 ° C./s or more is also an important production condition in the present invention. Due to the ferrite transformation from untransformed austenite during reheating and the subsequent discharge of C into untransformed austenite, untransformed austenite enriched with C during air cooling after reheating is transformed into MA (island martensite). To do. In order to obtain such MA (island martensite), it is necessary to reheat from the temperature above the Bf point to a temperature range of 550 to 750 ° C. after accelerated cooling.
- the rate of temperature increase is less than 0.5 ° C./s, it takes a long time to reach the target reheating temperature, so that the production efficiency deteriorates, and pearlite transformation occurs, so that MA (island martensite) is obtained. Therefore, a sufficiently low yield ratio cannot be obtained.
- the reheating temperature is less than 550 ° C., the ferrite transformation does not occur sufficiently and C is not sufficiently discharged into the untransformed austenite, so that MA is not generated and a low yield ratio cannot be achieved.
- the reheating temperature exceeds 750 ° C., the hardness difference between ferrite and bainite becomes less than 60 in HV due to softening of bainite, and a low yield ratio cannot be achieved.
- the reheating temperature range is set to 550 to 750 ° C.
- the cooling rate after reheating is basically preferably air cooling.
- the present invention is a steel pipe using the above-described steel material.
- Examples of the method for forming a steel pipe include a method for forming a steel pipe into a shape by cold forming such as a UOE process or a press bend (also called a bending press).
- the end bending of the width direction end of the steel plate is performed using a press machine, and then the steel plate is processed using a press machine.
- the steel plate is formed into a cylindrical shape so that the widthwise ends of the steel plate face each other.
- the opposing widthwise ends of the steel plates are brought together and welded. This welding is called seam welding.
- seam welding a cylindrical steel plate is constrained, the widthwise ends of opposing steel plates are butted against each other in a tack welding process, and welding is performed on the inner and outer surfaces of the butt portion of the steel plate by the submerged arc welding method.
- a method having a two-stage process that is, a main welding process for performing the above-described process is preferable.
- pipe expansion is performed to remove residual welding stress and improve roundness of the steel pipe.
- the pipe expansion ratio ratio of the outer diameter change amount before and after the pipe expansion to the outer diameter of the pipe before the pipe expansion
- the tube expansion rate is preferably in the range of 0.5% to 1.2%.
- a coating treatment can be carried out for the purpose of preventing corrosion.
- a known resin may be applied to the outer surface after heating to a temperature range of 200 to 300 ° C., for example.
- a steel pipe having a substantially circular cross-sectional shape is manufactured by successively forming a steel plate by repeating three-point bending. Thereafter, seam welding is performed in the same manner as the above-described UOE process. Also in the case of press bend, tube expansion may be performed after seam welding, and coating may also be performed.
- Steels having a thickness of 30 mm and a thickness of 33 mm were manufactured using steels (steel types A to K) having the composition shown in Table 1 (the balance being Fe and inevitable impurities) under the conditions shown in Table 2.
- the reheating was performed using an induction heating furnace or a gas combustion furnace.
- temperature such as heating temperature, rolling completion temperature, cooling stop (finishing) temperature, and reheating temperature, was made into the center part temperature of a steel plate.
- the center temperature was calculated by inserting a thermocouple in the center of the slab or steel plate and directly measuring it, or calculating the surface temperature of the slab or steel plate using parameters such as plate thickness and thermal conductivity.
- the cooling rate is an average cooling rate obtained by dividing the temperature difference required for cooling to the cooling stop (end) temperature after the hot rolling is finished by the time required for the cooling.
- the reheating rate (temperature increase rate) is an average temperature increase rate divided by the time required to reheat the temperature difference necessary for reheating up to the reheating temperature after cooling.
- the steel material produced as described above was subjected to structure observation and measured for tensile properties, hardness difference, and HIC resistance.
- the evaluation method is as follows.
- (1) Microstructural observation A specimen for microstructural observation is collected from the obtained thick steel plate, the L-direction cross section is polished, the nital corrosion is performed, and an optical microscope is used for the central portion of the plate thickness that is ⁇ 2 mm from the central plate thickness position. (Magnification: 400 times) or scanning electron microscope (magnification: 2000 times), the microstructure is observed for three or more fields of view, imaged, and image analysis is performed to determine the type of tissue and the area fraction of each phase. It was. (2) Tensile properties About the tensile strength before strain aging treatment, two No.
- HIC resistance is HIC tested for 96 hours in a 1 mol / l acetate buffer solution containing 5% NaCl with a pH of about 5.0 saturated with 100% hydrogen sulfide. The case where it was not recognized was judged as having good anti-HIC characteristics and indicated by a circle, and the case where a crack occurred was indicated by an x.
- Table 3 shows the measurement results. In addition, except for what has been specially stated (No. 11 is the area fraction of ferrite: 2%), the area fraction of ferrite is 10% or more, and the area ratio of bainite is 10% or more. there were.
- the structure of the steel material is composed of ferrite, bainite and island martensite.
- the area fraction of island martensite is 0.5-5%, and the hardness difference between ferrite and bainite is 60 or more in terms of Vickers hardness. Met.
- the chemical composition is within the scope of the present invention, but the production method is outside the scope of the present invention. Therefore, any of the structure, strength, yield ratio before and after strain aging treatment, and uniform elongation can be selected. It was insufficient.
- No. 14 to 18, the chemical components are outside the scope of the present invention or the production method is outside the scope of the present invention, so that sufficient strength cannot be obtained, the yield ratio is high, the uniform elongation is low, the HIC test Cracking occurred.
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Abstract
Description
(a)Cuを適量添加し、Moを含有しない、あるいは、含有しても0.01%以下にすることにより、耐HIC特性を向上させることが可能である。
(b)加速冷却過程で、冷却開始温度を適正に制御し、ベイナイト変態途中、すなわち未変態オーステナイトが存在する温度領域で冷却を停止し、その後ベイナイト変態終了温度(以下、Bf点と記載する。)以上から再加熱を行うことにより、鋼板の金属組織を、フェライトおよびベイナイトの混合相中に硬質相である島状マルテンサイト(Martensite-Austenite constituent、以下、MAと記載する。)が均一に生成した3相組織となり、歪時効処理前および歪時効処理後(以下、「歪時効処理前後」と称することもある。)の低降伏比化が可能である。
(c)加速冷却における冷却開始温度と冷却停止温度を適正な温度とすることで、固溶Cを低減することができるため、歪時効後の降伏比の上昇が抑制できる。
[1]成分組成として、質量%で、C:0.030~0.100%、Si:0.01~0.50%、Mn:0.5~2.5%、P:0.015%以下、S:0.002%以下、Cu:0.20~1.00%、Mo:0.01%以下、Nb:0.005~0.05%、Ti:0.005~0.040%、Al:0.10%以下、N:0.007%以下を含有し、残部Feおよび不可避的不純物からなる成分組成を有し、金属組織がフェライトとベイナイトと島状マルテンサイトとを有し、前記島状マルテンサイトの面積分率が0.5~5.0%であり、前記フェライトと前記ベイナイトとの硬度差がビッカース硬さで60以上であり、300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下である耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材。
[2]前記成分組成に、さらに、質量%で、Ni:0.02~0.50%、Cr:1.00%以下、V:0.10%以下、Ca:0.0050%以下、B:0.0050%以下の1種または2種以上を含有する[1]に記載の耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材。
[3][1]または[2]に記載の成分組成を有する鋼を、1000~1300℃の温度に加熱し、Ar3点以上の圧延終了温度で熱間圧延した後、(Ar3-50)~(Ar3+30)℃の冷却開始温度から5℃/s以上の冷却速度で冷却停止温度450~650℃まで加速冷却を行い、その後直ちに0.5℃/s以上の昇温速度で550℃~750℃まで再加熱を行う、金属組織がフェライトとベイナイトと島状マルテンサイトとを有し、前記島状マルテンサイトの面積分率が0.5~5%であり、前記フェライトと前記ベイナイトの硬度差がビッカース硬さで60以上であり、300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下である耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材の製造方法。
[4][1]または[2]に記載の鋼材を素材とする溶接鋼管。
なお、本発明における耐歪時効特性とは、300℃以下の温度の熱処理を施しても降伏比の過度な上昇を抑制できる特性をいう。また、本発明における耐HIC特性とは、pH5以上の湿潤硫化水素環境において水素誘起割れが発生しない特性をいう。また、高変形能とは、一様伸びが9%以上および降伏比が90%以下を満たす特性をいう。
以下に、本発明に係る鋼材の成分組成の限定理由を説明する。なお、成分組成を示す単位の%は、全て質量%を意味する。
Cは炭化物として析出強化に寄与する元素である。Cが0.030%未満では、MA(島状マルテンサイト)の生成が不足するため、十分な強度が確保できないほか、フェライトとベイナイトとの硬度差を所定量確保できないため、降伏比が大きくなる。Cが0.100%を超えると、靭性や溶接性の劣化、歪時効による降伏比の上昇を招く。このため、C含有量を0.030~0.100%に規定する。好ましくは、C含有量は0.05%以上である。また、好ましくは、C含有量は0.09%以下である。
Siは脱酸のため含有する。Siが0.01%未満では脱酸効果が十分ではない。Siが0.50%を超えると靭性や溶接性の劣化を招く。このため、Si含有量を0.01~0.50%に規定する。好ましくは、Si含有量は0.01~0.3%である。
Mnは、強度、靭性のため含有する。Mnが0.5%未満ではその効果が十分ではなく、また、MA(島状マルテンサイト)が不足するため、降伏比が大きくなる。このため、Mn含有量は0.5%以上とし、MA生成による低降伏比化の観点から、好ましくは、1.2%以上であり、より好ましくは、1.5%以上である。一方、Mnが2.5%を超えると靭性と溶接性が劣化する。このため、Mn含有量を2.5%以下に規定し、好ましくは、2.0%以下である。
Pは溶接性と耐HIC特性を劣化させる不可避的不純物元素である。このため、P含有量は、0.015%以下に規定する。好ましくは、P含有量は0.010%以下である。
Sは一般的には鋼中においてはMnS介在物となり耐HIC特性を劣化させる。このため、少ないほどよい。Sが0.002%以下であれば問題ないため、S含有量の上限を0.002%に規定する。好ましくは、S含有量は0.0015%以下である。
Cuは本発明において重要な元素であり、鋼中への水素の侵入を抑制し、耐HIC特性向上に寄与する。しかし、Cuが0.20%未満ではその効果が十分ではなく、1.00%を超えると溶接性が劣化する。このため、Cu含有量を0.20~1.00%に規定する。好ましくは、Cu含有量は0.25%以上である。また、好ましくは、Cu含有量は0.5%以下である。
Moは歪時効による降伏比の上昇、および、耐HIC特性の劣化を招く。このため、Moは含有しないか、あるいは含有しても0.01%以下に規定する。好ましくは、Mo含有量は0.005%以下である。
Nbは組織の微細化により靭性を向上させ、さらに炭化物を形成し、強度上昇に寄与する。しかし、Nbが0.005%未満ではその効果が十分ではなく、0.05%を超えると溶接熱影響部の靭性が劣化する。このため、Nb含有量を0.005~0.05%に規定する。好ましくは、Nb含有量は0.01~0.05%である。
TiはTiNのピニング効果により、スラブ加熱時のオーステナイト粗大化を抑制し、母材靭性を向上させ、さらに固溶Nを低減し歪時効による降伏比上昇を抑制する。しかし、Tiが0.005%未満ではその効果が十分ではなく、0.040%を超えると溶接熱影響部の靭性が劣化する。このため、Ti含有量は0.005~0.040%に規定する。好ましくは、Ti含有量は0.005~0.02%である。
Alは脱酸剤として含有する。Alが0.10%を超えると鋼の清浄度が低下し、靭性が劣化する。このため、Al含有量は0.10%以下に規定する。好ましくは、Al含有量は0.01~0.08%とする。
Nは歪時効による降伏比の上昇、溶接熱影響部の靭性の劣化を招く不可避的不純物元素である。このため、N含有量の上限を0.007%に規定する。好ましくは、N含有量は0.006%以下である。
Niは耐HIC向上に寄与し、靭性の改善と強度の上昇に有効な元素である。Niが0.02%未満ではその効果が十分ではなく、0.50%を超えて含有しても効果が飽和し、むしろコスト的に不利になる。このため、含有する場合はNi含有量を0.02~0.50%に規定する。好ましくは、Ni含有量は0.2%以上である。また、好ましくは、Ni含有量は0.4%以下である。
Crは低Cでも十分な強度を得るために有効な元素である。Crが1.00%を超えると溶接性が劣化する。このため、含有する場合はCr含有量の上限を1.00%に規定する。好ましくは、0.1~0.5%以下である。
Vは組織の微細化により靭性を向上させ、さらに炭化物を形成し、強度の向上に寄与する。Vが0.10%を超えると溶接熱影響部の靭性が劣化する。このため、含有する場合はV含有量は0.10%以下に規定する。好ましくは、V含有量は0.005%以上である。また、好ましくは、V含有量は0.05%以下である。
Caは硫化物系介在物の形態制御による靭性改善に有効な元素である。Caが0.0050%を超えると効果が飽和し、むしろ、鋼の清浄度の低下により靭性を劣化させる。このため、含有する場合は、Ca含有量を0.0050%以下に規定する。好ましくは、Ca含有量は0.001%以上である。また、好ましくは、Ca含有量は0.004%以下である。
Bは強度上昇、溶接熱影響部の靭性改善に有効な元素である。Bが0.0050%を超えると溶接性を劣化させる。このため、含有する場合は、B含有量を0.0050%以下に規定する。好ましくは、B含有量は0.003%以下である。また、好ましくは、B含有量は0.0003%以上である。
本発明の鋼板の金属組織は、フェライトとベイナイトと島状マルテンサイトとからなる3相組織を主体とする。フェライトとベイナイトと島状マルテンサイトとからなる3相組織を主体とするとは、フェライトとベイナイトと島状マルテンサイトの面積分率が合計で90%以上の複相組織をいう。残部としては、マルテンサイト(島状マルテンサイトを除く)やパーライト、残留オーステナイト等から選ばれる1種または2種以上の合計の面積分率が10%以下の組織である。
300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下
地震地帯に適用されるラインパイプ用鋼材は、地盤変動のような大きな変形を受ける場合でも破壊しないように高変形能であることが要求されている。さらに防食のためのコーティングで最大300℃に加熱される歪時効処理後でも高変形能を維持することが必要である。300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下である場合は、十分な高変形能が得られ、地震などの大変形により破壊に至る虞はない。そのため、本発明の鋼材では、300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びを9%以上とし、降伏比を90%以下とする。この300℃以下の温度の歪時効処理前および歪時効処理後の夫々は、高変形能の観点から、一様伸びが10%以上および降伏比が88%以下であることが好ましい。
次に、本発明の高変形能ラインパイプ用鋼材の製造方法について説明する。本発明の高変形能ラインパイプ用鋼材の製造方法としては、上述した成分組成を有する鋼素材を用い、加熱温度:1000~1300℃、圧延終了温度:Ar3点以上で熱間圧延を行った後、(Ar3-50)~(Ar3+30)℃の冷却開始温度から5℃/s以上の冷却速度で冷却停止温度450~650℃まで加速冷却を行い、その後直ちに0.5℃/s以上の昇温速度で550~750℃まで再加熱を行うことで、所望の金属組織とすることができる。ここで、温度は鋼材の中央部温度とする。なお、Ar3点は、以下の式より計算される。
Ar3(℃)=910-310C-80Mn-20Cu-15Cr-55Ni-80Mo
上記式において、元素記号は各元素の含有量(質量%)を示し、含有しない場合は0とする。
加熱温度が1000℃未満では炭化物の固溶が不十分で必要な強度が得られず、1300℃を超えると母材靭性が劣化する。このため、加熱温度を1000~1300℃に規定する。
圧延終了温度がAr3点未満であると、その後のフェライト変態速度が低下し、圧延による塑性歪がフェライト中に残存してフェライト強度が高くなり、フェライトとベイナイトとの硬度差が低下するため、所望の降伏比が達成できなくなる。このため、圧延終了温度をAr3点以上に規定する。さらに、900℃以下の温度域における累積圧下率を50%以上とすることが好ましい。900℃以下の温度域における累積圧下率を50%以上とすることにより、オーステナイト粒を微細化することができる。
冷却開始温度が(Ar3-50)℃未満の温度ではフェライトの面積分率が増加し、母材強度が劣化する。さらに、フェライトとベイナイトの硬度差が大きくなり、耐HIC特性が劣化する。よって、冷却開始温度は(Ar3-50)℃以上とし、好ましくは、(Ar3-30)℃以上である。また、冷却開始温度が(Ar3+30)℃を超えるとフェライトの面積分率が減少するとともにフェライトとベイナイトの硬度差の低下が生じ、低降伏比化を達成するには不十分となる。よって、冷却開始温度は(Ar3+30)℃以下とし、好ましくは、(Ar3+25)℃以下である。
冷却速度が5℃/s未満では冷却時にパーライトを生成し、十分な強度や低降伏比が得られないため、冷却速度を5℃/s以上に規定する。好ましくは、冷却速度は8℃/s以上、より好ましくは10℃/s以上である。また、好ましくは、冷却速度は100℃/s以下、より好ましくは60℃/s以下である。
本発明において、加速冷却の冷却停止温度は重要な製造条件である。本発明では、再加熱後に存在するCの濃縮した未変態オーステナイトが、その後の空冷時にMA(島状マルテンサイト)へと変態する。すなわち、ベイナイト変態途中の未変態オーステナイトが存在する温度域で冷却を停止する必要がある。冷却停止温度が450℃未満では、ベイナイト変態が完了するため空冷時にMA(島状マルテンサイト)が生成せず低降伏比化が達成できない。650℃を超えると冷却中に析出するパーライトにCが消費されMA(島状マルテンサイト)の生成が抑制され、MA量が不足する。このため、加速冷却の冷却停止温度を450~650℃に規定する。好ましくは、冷却停止温度は515℃以上、より好ましくは530℃以上である。また、好ましくは、冷却停止温度は635℃以下、より好ましくは620℃以下である。
このプロセスも、本発明において重要な製造条件である。再加熱時の未変態オーステナイトからのフェライト変態と、それに伴う未変態オーステナイトへのCの排出により、再加熱後の空冷時にCが濃化した未変態オーステナイトがMA(島状マルテンサイト)へと変態する。このようなMA(島状マルテンサイト)を得るためには、加速冷却後Bf点以上の温度から550~750℃の温度域まで再加熱する必要がある。昇温速度が0.5℃/s未満では、目的の再加熱温度に達するまでに長時間を要するため製造効率が悪化し、また、パーライト変態が生じるためMA(島状マルテンサイト)が得られず、十分な低降伏比を得ることができない。
再加熱の温度が550℃未満ではフェライト変態が十分起こらずCの未変態オーステナイトへの排出が不十分となり、MAが生成せず低降伏比化が達成できない。再加熱温度750℃を超えると、ベイナイトの軟化によりフェライトとベイナイトとの硬度差がHVで60未満となり、低降伏比が達成できない。このため、再加熱の温度域を550~750℃に規定する。なお、確実にフェライト変態を生じさせてCを未変態オーステナイトへ濃化させるためには、再加熱の際に、再加熱開始温度より50℃以上昇温することが望ましい。再加熱後の冷却速度は基本的には空冷とすることが好ましい。
さらに、溶接鋼管の製造方法について説明する。
(1)組織観察
得られた厚鋼板から組織観察用試験片を採取し、L方向断面を研磨、ナイタール腐食して、板厚中央位置から±2mmの領域である板厚中央部について、光学顕微鏡(倍率:400倍)または走査型電子顕微鏡(倍率:2000倍)を用いて、ミクロ組織を各3視野以上観察し、撮像して画像解析により、組織の種類および各相の面積分率を求めた。
(2)引張特性
歪時効処理前の引張強度については、圧延垂直方向のJIS Z 2201に規定される4号試験片を2本採取し、引張試験を行い、その平均値で評価した。引張強度517MPa以上(API 5L X60以上)を本発明に必要な強度とした。降伏比、一様伸びは、圧延方向のJIS Z 2201に規定される4号試験片を2本採取し、引張試験を行い、その平均値で評価した。降伏比90%以下、一様伸び9%以上を本発明に必要な降伏比とした。
また、歪時効処理後の引張強度については、圧延方向のJIS Z 2201に規定される4号試験片を2本採取し、2.0%の引張歪を付与した後、250℃にて5分間保持して、歪時効処理した後、引張試験を実施し、その平均値で評価した。なお、歪時効処理後の評価基準は、上述した歪時効処理前の評価基準と同一の基準で判定した。
(3)硬度差
得られた厚鋼板から硬さ測定用試験片を採取し、フェライトとベイナイトの硬度を、測定荷重5gのビッカース硬度計により測定し、10点以上の測定結果の平均値を用いて、フェライトとベイナイトとの硬度差を求めた。
(4)耐HIC特性
耐HIC特性は100%硫化水素を飽和させたpH約5.0の5%NaClを含む1mol/l酢酸緩衝溶液中に96時間浸漬する条件でHIC試験を行い、割れが認められない場合を耐HIC特性良好と判断して○印で示し、割れが発生した場合を×印で示した。
Claims (4)
- 成分組成として、質量%で、C:0.030~0.100%、Si:0.01~0.50%、Mn:0.5~2.5%、P:0.015%以下、S:0.002%以下、Cu:0.20~1.00%、Mo:0.01%以下、Nb:0.005~0.05%、Ti:0.005~0.040%、Al:0.10%以下、N:0.007%以下を含有し、残部Feおよび不可避的不純物からなる成分組成を有し、金属組織がフェライトとベイナイトと島状マルテンサイトとを有し、前記島状マルテンサイトの面積分率が0.5~5.0%であり、前記フェライトと前記ベイナイトとの硬度差がビッカース硬さで60以上であり、300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下である耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材。
- 前記成分組成に、さらに、質量%で、Ni:0.02~0.50%、Cr:1.00%以下、V:0.10%以下、Ca:0.0050%以下、B:0.0050%以下の1種または2種以上を含有する請求項1に記載の耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材。
- 請求項1または請求項2に記載の成分組成を有する鋼を、1000~1300℃の温度に加熱し、Ar3点以上の圧延終了温度で熱間圧延した後、(Ar3-50)~(Ar3+30)℃の冷却開始温度から5℃/s以上の冷却速度で冷却停止温度450~650℃まで加速冷却を行い、その後直ちに0.5℃/s以上の昇温速度で550℃~750℃まで再加熱を行う、金属組織がフェライトとベイナイトと島状マルテンサイトとを有し、前記島状マルテンサイトの面積分率が0.5~5%であり、前記フェライトと前記ベイナイトの硬度差がビッカース硬さで60以上であり、300℃以下の温度の歪時効処理前および歪時効処理後の夫々について、一様伸びが9%以上および降伏比が90%以下である耐歪時効特性及び耐HIC特性に優れた高変形能ラインパイプ用鋼材の製造方法。
- 請求項1または請求項2に記載の鋼材を素材とする溶接鋼管。
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US10465261B2 (en) | 2019-11-05 |
US20170022583A1 (en) | 2017-01-26 |
RU2016138771A (ru) | 2018-04-02 |
JPWO2015151468A1 (ja) | 2017-04-13 |
KR101885234B1 (ko) | 2018-08-03 |
EP3128030A4 (en) | 2017-11-29 |
EP3128030A1 (en) | 2017-02-08 |
EP3128030B1 (en) | 2020-11-11 |
CN106164314A (zh) | 2016-11-23 |
KR20160129056A (ko) | 2016-11-08 |
JP6048615B2 (ja) | 2016-12-21 |
RU2016138771A3 (ja) | 2018-04-02 |
RU2653740C2 (ru) | 2018-05-14 |
CN106164314B (zh) | 2018-10-30 |
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