JP5621478B2 - High toughness and high deformation steel plate for high strength steel pipe and method for producing the same - Google Patents
High toughness and high deformation steel plate for high strength steel pipe and method for producing the same Download PDFInfo
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- 229910000831 Steel Inorganic materials 0.000 title claims description 97
- 239000010959 steel Substances 0.000 title claims description 97
- 238000004519 manufacturing process Methods 0.000 title claims description 20
- 238000001816 cooling Methods 0.000 claims description 53
- 229910000734 martensite Inorganic materials 0.000 claims description 38
- 229910001566 austenite Inorganic materials 0.000 claims description 33
- 238000000926 separation method Methods 0.000 claims description 30
- 229910000859 α-Fe Inorganic materials 0.000 claims description 28
- 238000005096 rolling process Methods 0.000 claims description 27
- 230000009466 transformation Effects 0.000 claims description 26
- 229910001563 bainite Inorganic materials 0.000 claims description 23
- 238000005098 hot rolling Methods 0.000 claims description 21
- 238000000034 method Methods 0.000 claims description 18
- 239000002184 metal Substances 0.000 claims description 17
- 229910052751 metal Inorganic materials 0.000 claims description 17
- 238000009863 impact test Methods 0.000 claims description 16
- 230000009467 reduction Effects 0.000 claims description 16
- 238000012360 testing method Methods 0.000 claims description 11
- 239000000203 mixture Substances 0.000 claims description 10
- 230000001186 cumulative effect Effects 0.000 claims description 9
- 239000012535 impurity Substances 0.000 claims description 7
- 229910052717 sulfur Inorganic materials 0.000 claims description 6
- 229910052791 calcium Inorganic materials 0.000 claims description 5
- 229910052757 nitrogen Inorganic materials 0.000 claims description 5
- 229910052748 manganese Inorganic materials 0.000 claims description 4
- 229910052760 oxygen Inorganic materials 0.000 claims description 4
- 229910052750 molybdenum Inorganic materials 0.000 claims description 3
- 229910052698 phosphorus Inorganic materials 0.000 claims description 3
- 229910052720 vanadium Inorganic materials 0.000 claims description 3
- 229910052799 carbon Inorganic materials 0.000 claims description 2
- 229910052804 chromium Inorganic materials 0.000 claims description 2
- 229910052759 nickel Inorganic materials 0.000 claims description 2
- 229910052758 niobium Inorganic materials 0.000 claims description 2
- 229910052742 iron Inorganic materials 0.000 claims 1
- 238000003303 reheating Methods 0.000 description 21
- 239000000463 material Substances 0.000 description 19
- 230000007423 decrease Effects 0.000 description 13
- 238000010438 heat treatment Methods 0.000 description 13
- 230000000694 effects Effects 0.000 description 12
- 230000000052 comparative effect Effects 0.000 description 11
- VNWKTOKETHGBQD-UHFFFAOYSA-N methane Chemical compound C VNWKTOKETHGBQD-UHFFFAOYSA-N 0.000 description 8
- 238000001953 recrystallisation Methods 0.000 description 8
- 238000010521 absorption reaction Methods 0.000 description 7
- 239000010953 base metal Substances 0.000 description 6
- 230000008569 process Effects 0.000 description 6
- 230000015572 biosynthetic process Effects 0.000 description 5
- 239000007789 gas Substances 0.000 description 5
- 238000005204 segregation Methods 0.000 description 5
- 239000006104 solid solution Substances 0.000 description 5
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 4
- 239000003345 natural gas Substances 0.000 description 4
- 238000009628 steelmaking Methods 0.000 description 4
- 238000003466 welding Methods 0.000 description 4
- 230000002411 adverse Effects 0.000 description 3
- 238000009749 continuous casting Methods 0.000 description 3
- 238000004881 precipitation hardening Methods 0.000 description 3
- 239000002436 steel type Substances 0.000 description 3
- 230000003749 cleanliness Effects 0.000 description 2
- 238000003776 cleavage reaction Methods 0.000 description 2
- 229910052802 copper Inorganic materials 0.000 description 2
- 238000005530 etching Methods 0.000 description 2
- 238000010191 image analysis Methods 0.000 description 2
- 230000001771 impaired effect Effects 0.000 description 2
- 230000006872 improvement Effects 0.000 description 2
- 239000011159 matrix material Substances 0.000 description 2
- 150000001247 metal acetylides Chemical class 0.000 description 2
- 238000010791 quenching Methods 0.000 description 2
- 230000000171 quenching effect Effects 0.000 description 2
- 238000011160 research Methods 0.000 description 2
- 229920006395 saturated elastomer Polymers 0.000 description 2
- 230000007017 scission Effects 0.000 description 2
- 239000000126 substance Substances 0.000 description 2
- 238000009864 tensile test Methods 0.000 description 2
- 229910001035 Soft ferrite Inorganic materials 0.000 description 1
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 1
- 238000004364 calculation method Methods 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 238000002485 combustion reaction Methods 0.000 description 1
- 239000012141 concentrate Substances 0.000 description 1
- 238000010276 construction Methods 0.000 description 1
- 239000010779 crude oil Substances 0.000 description 1
- 239000013078 crystal Substances 0.000 description 1
- 230000006866 deterioration Effects 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 230000006698 induction Effects 0.000 description 1
- 239000004615 ingredient Substances 0.000 description 1
- 238000011835 investigation Methods 0.000 description 1
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
- 230000007246 mechanism Effects 0.000 description 1
- 238000002844 melting Methods 0.000 description 1
- 230000008018 melting Effects 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 230000003287 optical effect Effects 0.000 description 1
- 229910001562 pearlite Inorganic materials 0.000 description 1
- 238000005498 polishing Methods 0.000 description 1
- 239000002244 precipitate Substances 0.000 description 1
- 230000002265 prevention Effects 0.000 description 1
- 238000012545 processing Methods 0.000 description 1
- 230000000717 retained effect Effects 0.000 description 1
- 238000005728 strengthening Methods 0.000 description 1
- 238000005496 tempering Methods 0.000 description 1
- 150000003568 thioethers Chemical class 0.000 description 1
- 230000007704 transition Effects 0.000 description 1
- 230000000007 visual effect Effects 0.000 description 1
Classifications
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
- B21B—ROLLING OF METAL
- B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
-
- 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
- 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
- C22C38/18—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
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Mechanical Engineering (AREA)
- Materials Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Heat Treatment Of Steel (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
Description
本発明は、高張力厚鋼板およびその製造方法に関し、引張強度６００ＭＰａを超える高強度と、降伏比が８０％以下の優れた変形能と、シャルピー衝撃試験やＤＷＴＴ（Ｄｒｏｐ Ｗｅｉｇｈｔ Ｔｅａｒ Ｔｅｓｔ：落重引裂試験）などで評価した場合の高靭性と、を兼ね備え、かつ、パイプとしたときの延性亀裂伝播停止性能にも優れる、天然ガス輸送用の鋼管素材として好適な鋼板およびその製造方法に関する。 The present invention relates to a high-strength thick steel plate and a method for producing the same, high strength exceeding a tensile strength of 600 MPa, excellent deformability with a yield ratio of 80% or less, Charpy impact test, DWTT (Drop Weight Tear Test: drop weight tearing) The present invention relates to a steel sheet suitable as a steel pipe material for transporting natural gas, which has high toughness when evaluated in a test) and has excellent ductility crack propagation stopping performance when used as a pipe, and a method for producing the same.
近年、天然ガスや原油の輸送用として使用されるラインパイプは、高圧化による輸送効率の向上や薄肉化による現地溶接施工効率の向上のため、年々高強度化している。特に、天然ガス輸送パイプラインにおいては、パイプに延性亀裂が発生した場合、亀裂部からのガス流出によるガス圧低下速度より、亀裂伝播速度の方が速い場合には、亀裂開口力が低下しないため、延々と亀裂が伝播してしまうことが想定され、亀裂伝播速度を抑えるために、ラインパイプ母材に高いシャルピー衝撃試験での吸収エネルギー（以下、単にシャルピー吸収エネルギーとも称する）を要求される。寒冷地にあるガス田からのガス輸送の場合には、−３０℃といった低温でも２００Jを超えるシャルピー吸収エネルギーを満足させなくてはならない場合がある。 In recent years, line pipes used for transportation of natural gas and crude oil have been strengthened year by year in order to improve transportation efficiency by increasing pressure and to improve local welding construction efficiency by reducing wall thickness. In particular, in a natural gas transportation pipeline, when a ductile crack occurs in the pipe, the crack opening force does not decrease when the crack propagation speed is faster than the gas pressure decrease speed due to gas outflow from the crack. In order to suppress the crack propagation rate, it is assumed that the crack propagates endlessly, and the line pipe base material is required to have high absorption energy in the Charpy impact test (hereinafter also simply referred to as Charpy absorption energy). In the case of gas transportation from a gas field in a cold region, it may be necessary to satisfy Charpy absorbed energy exceeding 200 J even at a low temperature of −30 ° C.
また、大地震や凍土地帯における地盤変動によって、ラインパイプに大変形が生じても、延性亀裂発生にいたらないための、高変形能の要求もなされるようになってきた。 In addition, there has been a demand for high deformability to prevent the occurrence of ductile cracks even if a large deformation occurs in a line pipe due to a large earthquake or ground deformation in a frozen land zone.
高変形能の指標として、降伏強度を引張強度で割った降伏比（ＹＲ）が使われ、低ＹＲ化されるほど亀裂発生の限界歪が向上する。 As an index of high deformability, the yield ratio (YR) obtained by dividing the yield strength by the tensile strength is used, and the lower the YR, the higher the limit strain for crack generation.
鋼材の金属組織を軟質なフェライトと、硬質なベイナイトやマルテンサイトなどが適度に分散した硬質相の２相組織とすることで、低ＹＲとなることが知られており、例えば特許文献１には、上記のような軟質相の中に硬質相が適度に分散した組織を得る製造方法として、焼入れ（Ｑ）と焼戻し（Ｔ）の中間に、フェライトとオーステナイト相の２相域からの焼入れ（Ｑ’）を施す熱処理方法が開示されている。 It is known that a low YR is obtained by making the metal structure of a steel material a two-phase structure of soft ferrite and a hard phase in which hard bainite and martensite are appropriately dispersed. As a production method for obtaining a structure in which a hard phase is moderately dispersed in the soft phase as described above, quenching from a two-phase region of ferrite and austenite phase (Q) between quenching (Q) and tempering (T) (Q A heat treatment method for applying ') is disclosed.
また、特許文献２には、軟質相を加工フェライトとしたフェライト＋ベイナイト＋マルテンサイト組織により低ＹＲ化が達成されることが開示されている。 Patent Document 2 discloses that a low YR can be achieved by a ferrite + bainite + martensite structure in which a soft phase is processed ferrite.
さらに、特許文献３には、ベイナイト中に島状マルテンサイトを分散させて低ＹＲと高シャルピー吸収エネルギーを両立させることが開示されている。 Further, Patent Document 3 discloses that island-like martensite is dispersed in bainite to achieve both low YR and high Charpy absorbed energy.
しかしながら、特許文献１に記載の技術では多数回の熱処理を行う必要があり、生産性が低下し、製造コストが上昇する。 However, in the technique described in Patent Document 1, it is necessary to perform heat treatment a number of times, which decreases productivity and increases manufacturing cost.
特許文献２に記載の技術による高靱性化は延性−脆性破面率の改善のために、フェライトの集合組織を積極的に発達させることによって得られるもので、シャルピー衝撃試験片の破面にはセパレーションが発生し、シャルピー吸収エネルギーはむしろ低下する。 High toughness by the technique described in Patent Document 2 is obtained by actively developing a texture of ferrite in order to improve the ductility-brittle fracture surface ratio. Separation occurs and the Charpy absorbed energy is rather lowered.
特許文献３に記載の技術による低ＹＲと高シャルピー吸収エネルギーの両立は、上述のフェライトの集合組織に起因したセパレーションの発生を抑制するよう、圧延終了温度をＡｒ３変態点を超える高い温度としているが、近年の研究では、このようなフェライトが生成しにくい圧延終了温度としても、セパレーション生成に伴うシャルピー吸収エネルギーの低下が起こりうることが判明した。
そこで、本発明は、生産性の低下や製造コストの上昇を伴わず、高強度、低ＹＲ、および高シャルピー吸収エネルギーを兼ね備えた厚鋼板およびその製造方法を提供することを目的とする。
The coexistence of low YR and high Charpy absorbed energy by the technique described in Patent Document 3 sets the rolling end temperature to a high temperature exceeding the Ar 3 transformation point so as to suppress the occurrence of separation due to the ferrite texture described above. However, recent research has revealed that the Charpy absorbed energy can be reduced due to the generation of separation even at the rolling end temperature at which such ferrite is difficult to generate.
Accordingly, an object of the present invention is to provide a thick steel plate having high strength, low YR, and high Charpy absorbed energy, and a method for manufacturing the same without lowering productivity and increasing manufacturing cost.
発明者等はまず、目標とする脆性亀裂伝播停止性能を得るために、許容されるセパレーション発生量を定量化した。図１に示すようなシャルピー衝撃試験片の破面上に生成するセパレーションに対し、図２に示すように、長さが1ｍｍ以上の全てのセパレーションの長さを測定し、それらの総和を破面断面積（すなわち、被検面の面積）で除した値（ＳＩ：セパレーションインデックス）を用い、シャルピー吸収エネルギーとの関係を求めたところ、図３に示すような関係が得られることがわかった。これらの結果より、試験温度が−３０℃のときに２００Jを超えるような高いシャルピー吸収エネルギーを得るためには、少なくともセパレーションインデックスを０．０５（ｍｍ−１）以下とする必要があることを知見した。
次に、ベイナイトを主体とする母相と、前記母相中に分散して存在する第２相として島状マルテンサイトを有する二相組織鋼板を製造するにあたり、特に、圧延終了温度をＡｒ３変態点を超える高い温度としても、シャルピー衝撃試験片にセパレーションが発生してしまう原因について鋭意研究を行った。
（１）その結果、延性-脆性遷移挙動を向上させるために実施する９５０℃以下のオーステナイト未再結晶温度域での強圧下により、オーステナイト組織に形成された集合組織が、その後の加速冷却の際に変態して生成するベイナイトに受け継がれること、
（２）および、この強圧下により、オーステナイト粒界での拡散変態が促進され、加速冷却を開始する温度がＡｒ３変態点以上であっても一部のオーステナイト粒界で初析フェライトが生成し、この粒界上のフェライトもまたセパレーション発生の原因となること、
（３）さらに、鋼中に存在するＭｎＳが、前記オーステナイト粒界からの初析フェライト生成を促進すること、ただしＣａ添加による硫化物形態制御によりＭｎＳをＣａＳのようなＣａ系介在物とすることで無害化できること、
を見出した。
The inventors first quantified the amount of separation that was allowed in order to obtain the desired brittle crack propagation stopping performance. For the separation generated on the fracture surface of the Charpy impact test piece as shown in FIG. 1, measure the length of all separations with a length of 1 mm or more as shown in FIG. Using the value (SI: separation index) divided by the cross-sectional area (that is, the area of the test surface) and determining the relationship with Charpy absorbed energy, it was found that the relationship shown in FIG. 3 was obtained. From these results, in order to obtain high Charpy absorbed energy exceeding 200 J when the test temperature is −30 ° C., it is found that at least the separation index needs to be 0.05 (mm −1 ) or less. did.
Next, in manufacturing a duplex steel sheet having a matrix mainly composed of bainite and island-like martensite as a second phase dispersed and present in the matrix, particularly the rolling end temperature is set to Ar 3 transformation. We conducted intensive research on the cause of separation in Charpy impact specimens even at high temperatures exceeding the point.
(1) As a result, the aggregate structure formed in the austenite structure is subjected to accelerated cooling in the subsequent abrupt pressure in the austenite non-recrystallization temperature range of 950 ° C. or less, which is performed to improve the ductile-brittle transition behavior. To be inherited by bainite produced by transformation to
(2) And under this strong pressure, diffusion transformation at the austenite grain boundary is promoted, and proeutectoid ferrite is generated at some austenite grain boundaries even when the temperature at which accelerated cooling starts is higher than the Ar 3 transformation point. The ferrite on this grain boundary can also cause separation.
(3) Further, MnS present in the steel promotes the formation of pro-eutectoid ferrite from the austenite grain boundaries, except that MnS is converted to a Ca-based inclusion such as CaS by controlling the form of sulfide by adding Ca. Can be detoxified,
I found.
本発明は上記した知見を元に、さらに検討を加えてなされたもので、その要旨は以下の通りである。 The present invention has been made on the basis of the above-described findings and further studies, and the gist thereof is as follows.
第一の発明は、質量％で、Ｃ：０．０４〜０．０８％、Ｓｉ：０．０５〜０．５％、Ｍｎ：１．８〜３．０％、Ｐ：０．００８％以下、Ｓ：０．０００６％以下、Ｏ：０．００３％以下、Ｎ：０．００６％以下、Ａｌ：０．００３〜０．０５％、Ｎｉ：０．１〜１．０％、Ｃｒ：０．０１〜０．５％、Ｎｂ：０．０１〜０．０５％、Ｔｉ：０．００５〜０．０２０％を含有し、残部Ｆｅおよび不可避的不純物からなり、金属組織がベイナイトを主体とし、前記ベイナイト中に第２相として島状マルテンサイトが面積率で５〜１５％均一分散し、旧オーステナイト粒界に存在するフェライト相の面積率が全金属組織の５％以下であって、−３０℃の試験温度でシャルピー衝撃試験を行った際の、シャルピー衝撃試験片の破面に存在するセパレーションが下記式（１）で定義されるセパレーションインデックスが０．０５（ｍｍ−１）以下であることを特徴とする、高靱性かつ高変形性高強度鋼管用鋼板である。 1st invention is the mass%, C: 0.04-0.08%, Si: 0.05-0.5%, Mn: 1.8-3.0%, P: 0.008% or less , S: 0.0006% or less, O: 0.003% or less, N: 0.006% or less, Al: 0.003-0.05%, Ni: 0.1-1.0%, Cr: 0 .01 to 0.5%, Nb: 0.01 to 0.05%, Ti: 0.005 to 0.020%, the balance being Fe and inevitable impurities, the metal structure being mainly bainite, In the bainite, island-like martensite as the second phase is uniformly dispersed in an area ratio of 5 to 15%, and the area ratio of the ferrite phase existing in the prior austenite grain boundaries is 5% or less of the total metal structure, and −30 Separators present on the fracture surface of Charpy impact test pieces when performing Charpy impact tests at a test temperature of ℃ And wherein the separation index ® emission is defined by the following formula (1) is 0.05 (mm -1) or less, a high toughness and high deformability strength steel pipe for a steel sheet.
第二の発明は、さらに、質量％で、Ｃｕ：０．１０〜１．０％、Ｍｏ：０．０１〜０．５％、Ｖ：０．００３〜０．０８％の中から選ばれる1種または２種以上を含有することを特徴とする第一の発明に記載の高靱性かつ高変形性高強度鋼管用鋼板である。 The second invention is further selected by mass% from Cu: 0.10 to 1.0%, Mo: 0.01 to 0.5%, V: 0.003 to 0.08% 1 A high toughness and high deformability high strength steel pipe steel plate according to the first invention, characterized by containing seeds or two or more.
第三の発明は、さらに、質量％で、Ｂ：０．０００５〜０．００３０％を含有し、かつ、ＢとＴｉとＮが下記式（２）を満足することを特徴とする第一または第二の発明に記載の高靱性かつ高変形性高強度鋼管用鋼板である。 The third invention is further characterized in that, in mass%, B: 0.0005 to 0.0030% is contained, and B, Ti, and N satisfy the following formula (2): It is a steel plate for high toughness and high deformability high strength steel pipes described in the second invention.
なお、上記式中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 In addition, [M] in the said formula shows the quantity (mass%) of the element M contained in steel.
第四の発明は、さらに、質量％で、Ｃａ：０．０００５〜０．００４０％を含有しかつ、ＣａとＯとＳが下記式（３）を満足することを特徴とする第一乃至第三の発明の何れかに記載の高靱性かつ高変形性高強度鋼管用鋼板である。
According to a fourth aspect of the present invention, the first to the second aspect is characterized by further containing, by mass%, Ca: 0.0005 to 0.0040%, and Ca, O, and S satisfy the following formula (3): A high-toughness and high-deformability high-strength steel pipe steel plate according to any one of the three inventions.
なお、上記式中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。
第五の発明は、第一乃至第四の発明の何れかに記載の組成を有する鋼を、連続鋳造して鋼片を製造し、該鋼片をＡｃ３変態点以上１１００℃以下に加熱後、熱間圧延を開始し、９５０℃以下の温度域における累積圧下率を４０〜６６％とする熱間圧延を行った後、Ａｒ３変態点以上の温度から冷却速度１０〜８０℃／ｓで加速冷却を開始し、４００〜６００℃の温度域で加速冷却を停止した後、ただちに、６００〜６８０℃の温度範囲に再加熱し、以後、空冷することを特徴とする高靱性かつ高変形性高強度鋼管用鋼板の製造方法である。
In addition, [M] in the said formula shows the quantity (mass%) of the element M contained in steel.
According to a fifth invention, a steel having the composition according to any one of the first to fourth inventions is continuously cast to produce a steel slab, and the steel slab is heated to an Ac 3 transformation point or higher and 1100 ° C. or lower. Then, after hot rolling was started and hot rolling was performed with a cumulative reduction in the temperature range of 950 ° C. or lower being 40 to 66%, the cooling rate was 10 to 80 ° C./s from the temperature above the Ar 3 transformation point. Highly tough and highly deformable, characterized by starting accelerated cooling, stopping accelerated cooling in the temperature range of 400 to 600 ° C., immediately reheating to 600 to 680 ° C., and then air cooling. It is a manufacturing method of the steel plate for high strength steel pipes.
第六の発明は、さらに、９５０℃以下の温度域での熱間圧延における各圧延パス１回当りの圧下率の平均が１０％以下であることを特徴とする第五の発明に記載の高靱性かつ高変形性高強度鋼管用鋼板の製造方法である。 According to a sixth aspect of the present invention, the average rolling reduction per rolling pass in the hot rolling at a temperature range of 950 ° C. or lower is 10% or lower. This is a method for producing a tough and highly deformable high strength steel pipe steel sheet.
本発明によれば、８０％以下の低降伏比と試験温度：−２０℃において８５％以上のＤＷＴＴ延性破面率、および試験温度：−３０℃において２００J以上のシャルピー衝撃値を達成する、高靱性かつ高変形性能を有する、天然ガス輸送用ラインパイプ用引張強度６００ＭＰａ超えの厚肉高張力鋼板の製造が可能となり、産業上極めて有用である。 According to the present invention, a low yield ratio of 80% or less and a test temperature: DWTT ductile fracture ratio of 85% or more at −20 ° C., and a Charpy impact value of 200 J or more at −30 ° C., high It is possible to produce a thick high-tensile steel sheet having a tensile strength exceeding 600 MPa for a natural gas transportation line pipe having toughness and high deformation performance, which is extremely useful industrially.
以下に本発明の成分組成、金属組織および製造方法について説明する。 The component composition, metal structure and production method of the present invention will be described below.
［成分組成］
本発明の鋼の成分組成を規定した理由を説明する。なお、成分％は全て質量％を意味する。
[Ingredient composition]
The reason for defining the composition of the steel of the present invention will be described. In addition, all component% means the mass%.
Ｃ：０．０４〜０．０８％
Ｃは本発明においては、鋼のベイナイト組織中に島状マルテンサイトを分散させ、低ＹＲを得るために重要な元素である。すなわち、ベイナイト組織より硬質な島状マルテンサイト中にＣを濃化させて生成させる必要があり、かつ、低ＹＲとするために十分な島状マルテンサイトの面積率を得るためには０．０４％以上の含有が必要であるが、０．０８％を超えて含有すると、板厚中央偏析部の島状マルテンサイトが増加し、シャルピー吸収エネルギーの低下を引き起こすため、Ｃ量は０．０４〜０．０８％の範囲とする。
C: 0.04 to 0.08%
In the present invention, C is an important element for obtaining a low YR by dispersing island martensite in the bainite structure of steel. That is, it is necessary to concentrate and form C in island martensite harder than the bainite structure, and 0.04 in order to obtain a sufficient area ratio of island martensite for a low YR. However, if the content exceeds 0.08%, island martensite in the central segregation portion of the plate thickness increases, causing a decrease in Charpy absorbed energy. The range is 0.08%.
Ｓｉ：０．０５〜０．５％
Ｓｉは固溶強化能を有する元素であり、０．０５％以上含有すると母材およびＨＡＺの強度を上昇させるので有効である。しかし、０．５％を超えて含有すると母材およびＨＡＺにおいて島状マルテンサイトが生成しやすくなる。特に、板厚中央偏析部のような、Ｍｎ、Ｐが濃化した領域でこの効果は顕著であり、母材シャルピー吸収エネルギーの低下を引き起こすため、Ｓｉ量は０．０５〜０．５％の範囲とする。
Si: 0.05-0.5%
Si is an element having a solid solution strengthening ability, and if contained at 0.05% or more, it is effective because it increases the strength of the base material and the HAZ. However, if the content exceeds 0.5%, island martensite is likely to be generated in the base material and HAZ. In particular, this effect is remarkable in a region where Mn and P are concentrated, such as a plate thickness central segregation portion, and causes a decrease in the base metal Charpy absorption energy. Therefore, the Si amount is 0.05 to 0.5%. Range.
Ｍｎ：１．８〜３．０％
Ｍｎは焼入性向上元素として作用する。さらに、多量に添加することで、フェライトに固溶できるＣ量を低減する効果があり、鋼のオーステナイト域から加速冷却でベイナイト変態させる際、未変態オーステナイト領域へのＣ濃化を大きくするので、島状マルテンサイトの生成量を増加させることができる。
後述のように、島状マルテンサイトの面積率を５％以上とするためには、少なくとも１．８％のＭｎ含有が必要である。一方、連続鋳造プロセスでは中心偏析部の濃度上昇が著しく、３．０％を超えて含有すると、多量な島状マルテンサイトの生成により母材吸収エネルギーの低下を招くため、Ｍｎ量は１．８〜３．０％の範囲とする。
Mn: 1.8-3.0%
Mn acts as a hardenability improving element. Furthermore, by adding a large amount, there is an effect of reducing the amount of C that can be dissolved in the ferrite, when increasing the concentration of C to the untransformed austenite region when bainite transformation by accelerated cooling from the austenite region of steel, The generation amount of island martensite can be increased.
As will be described later, in order to make the area ratio of the island-like martensite 5% or more, it is necessary to contain at least 1.8% of Mn. On the other hand, in the continuous casting process, the concentration of the central segregation part is remarkably increased, and if the content exceeds 3.0%, a large amount of island-like martensite is caused to decrease the base material absorption energy. The range is -3.0%.
Ｐ：０．００８％以下
Ｐは鋼中に不可避不純物として存在する。特に板厚中心部での偏析が著しい元素であり、島状マルテンサイトの増加を引き起こし、母材のシャルピー吸収エネルギーを低下させるため、Ｐ量の上限は０．００８％とする。好ましくは、０．００６％以下である。
P: 0.008% or less P is present as an inevitable impurity in steel. In particular, the segregation at the central portion of the plate thickness is an element that causes remarkable increase in island martensite and lowers the Charpy absorbed energy of the base material. Therefore, the upper limit of the P content is set to 0.008%. Preferably, it is 0.006% or less.
Ｓ：０．０００６％以下
Ｓもまた鋼中に不可避不純物として存在する。特に介在物として存在し、鋼の清浄度を低下させ、母材シャルピー吸収エネルギーに悪影響を及ぼすため、Ｓ量の上限は０．０００６％とする。好ましくは、０．０００４％以下である。
S: 0.0006% or less S is also present as an inevitable impurity in the steel. In particular, the upper limit of the amount of S is set to 0.0006% because it exists as inclusions and lowers the cleanliness of steel and adversely affects the base metal Charpy absorbed energy. Preferably, it is 0.0004% or less.
Ｏ：０．００３％以下
Ｏは、通常、鋼中の不可避不純物として存在し、酸化物系介在物の生成の原因となる。特に、０．００３％以上存在すると、粗大な介在物が生成し、母材のシャルピー吸収エネルギー低下を低下させるため、Ｏ量の上限は０．００３％とする。
O: 0.003% or less O is usually present as an inevitable impurity in steel and causes generation of oxide inclusions. In particular, when 0.003% or more is present, coarse inclusions are generated and the reduction in Charpy absorbed energy of the base material is reduced, so the upper limit of the O amount is set to 0.003%.
Ｎ：０．００６％以下
Ｎは、通常、鋼中の不可避不純物として存在するが、後述の通りＴｉ添加を行うことで、ＴｉＮを形成し、オーステナイト粒の粗大化を抑制する。ただし、Ｎ量が０．００６％を超える場合、溶接部、特に溶融線近傍の１４５０℃以上に加熱された領域でＴｉＮが分解すると固溶Ｎが鋼の靱性を著しく低下させるため、Ｎ量の上限は０．００６％とする。
N: 0.006% or less N is usually present as an unavoidable impurity in steel, but by adding Ti as described later, TiN is formed and coarsening of austenite grains is suppressed. However, if the amount of N exceeds 0.006%, if TiN decomposes in the welded portion, particularly in the region heated to 1450 ° C. or more near the melting line, solid solution N significantly reduces the toughness of the steel. The upper limit is 0.006%.
Ａｌ：０．００３〜０．０５％
Ａｌは脱酸元素として作用する。０．００３％以上の含有で十分な脱酸効果が得られるが、０．０５％超えて含有すると偏析部も含めて鋼の清浄度が低下し、靭性低下の原因となるため、Ａｌ量は０．００３〜０．０５％の範囲とする。
Al: 0.003-0.05%
Al acts as a deoxidizing element. A sufficient deoxidation effect can be obtained with a content of 0.003% or more, but if it exceeds 0.05%, the cleanliness of the steel including the segregation part is lowered and the toughness is reduced. The range is 0.003 to 0.05%.
Ｎｉ：０．１〜１．０％
Ｎｉは、焼入性向上元素として作用するほか、多量に添加しても靱性劣化を起こさないため、有用な元素である。この効果を得るために、０．１％以上の含有が必要であるが、高価な元素であるため、Ｎｉ量の上限は１．０％とする。
Ni: 0.1 to 1.0%
Ni is a useful element because it acts as a hardenability improving element and does not cause toughness deterioration even when added in a large amount. In order to acquire this effect, 0.1% or more must be contained. However, since it is an expensive element, the upper limit of the Ni content is 1.0%.
Ｃｒ：０．０１〜０．５％
Ｃｒもまた０．０１％以上の含有によって焼入性向上元素として作用し、多量のＭｎ添加の代替とすることができる。しかし、０．５％を超えて含有するとＨＡＺ靱性が著しく劣化するため、Ｃｒ量は０．０１〜０．５％の範囲とする。
Cr: 0.01 to 0.5%
Cr also acts as a hardenability improving element when contained in an amount of 0.01% or more, and can be used as a substitute for adding a large amount of Mn. However, if the content exceeds 0.5%, the HAZ toughness deteriorates remarkably, so the Cr content is in the range of 0.01 to 0.5%.
Ｎｂ：０．０１〜０．０５％
Ｎｂは炭化物を形成することで、特に２回以上の溶接熱サイクルを受ける溶接熱影響部（ＨＡＺ）の焼戻し軟化を防止して、引張強度６００ＭＰａを超える高強度ラインパイプ用鋼板として必要なＨＡＺ強度を得るために必要な元素である。
Nb: 0.01 to 0.05%
Nb forms carbides to prevent temper softening of the weld heat affected zone (HAZ) that is subjected to two or more welding heat cycles, and has a HAZ strength necessary as a steel plate for high-strength line pipes with a tensile strength exceeding 600 MPa. It is an element necessary for obtaining.
また、熱間圧延時のオーステナイト未再結晶温度領域を高温側に拡大する効果もあり、特に９５０℃まで未再結晶温度領域とするためには０．０１％以上の含有が必要である。一方、０．０５％を超えて含有すると、ベイナイト中の島状マルテンサイトの生成が著しくなり、母材のシャルピー吸収エネルギーを低下させることから、Ｎｂ量は０．０１〜０．０５％の範囲とする。 Moreover, there is also an effect of expanding the austenite non-recrystallization temperature region at the time of hot rolling to a high temperature side, and in order to make the non-recrystallization temperature region particularly up to 950 ° C., the content of 0.01% or more is necessary. On the other hand, if it exceeds 0.05%, the formation of island martensite in bainite becomes remarkable and the Charpy absorbed energy of the base material is reduced, so the Nb content is in the range of 0.01 to 0.05%. .
Ｔｉ：０．００５〜０．０２０％
Ｔｉは窒化物を形成し、鋼中の固溶Ｎ量低減に有効であるほか、析出したＴｉＮがピンニング効果でオーステナイト粒の粗大化を抑制することで、母材およびＨＡＺの靱性向上に寄与する。必要なピンニング効果を得るためには０．００５％以上の含有が必要であるが、０．０２０％を超えて含有すると炭化物を形成するようになり、その析出硬化で靱性が著しく劣化するため、Ｔｉ量は０．００５〜０．０２０％の範囲とする。
Ti: 0.005-0.020%
Ti forms nitrides and is effective in reducing the amount of solute N in the steel, and the precipitated TiN suppresses the austenite grain coarsening by the pinning effect, thereby contributing to the improvement of the toughness of the base material and the HAZ. . In order to obtain the required pinning effect, 0.005% or more is necessary, but if it exceeds 0.020%, carbides are formed, and the toughness is significantly deteriorated due to precipitation hardening. The Ti amount is in the range of 0.005 to 0.020%.
本発明の基本成分組成は以上であるが、更に強度を上昇させる場合、Ｃｕ、Ｍｏ、Ｖの１種または２種以上を添加することができる。また、靭性向上の観点から、Ｂ、Ｃａを選択元素として含有することもできる。 Although the basic component composition of the present invention is as described above, one or more of Cu, Mo, and V can be added to further increase the strength. Moreover, B and Ca can also be contained as a selective element from a viewpoint of a toughness improvement.
Ｃｕ：０．１〜１．０％
Ｃｕは０．１％以上の含有によって焼入性向上元素として作用し、多量のＭｎ添加の代替とすることができる。しかし、１．０％を超えて含有すると、過飽和に固溶したＣｕが加速冷却後の再加熱時に析出し、特に鋼の降伏強度が析出硬化によって上昇する結果、低ＹＲとすることが困難となるため、Ｃｕを含有する場合は、０．１〜１．０％の範囲とすることが好ましい。
Cu: 0.1 to 1.0%
When Cu is contained in an amount of 0.1% or more, it acts as a hardenability improving element and can be used as a substitute for adding a large amount of Mn. However, if it exceeds 1.0%, Cu dissolved in supersaturation is precipitated during reheating after accelerated cooling, and the yield strength of steel is increased by precipitation hardening, and as a result, it is difficult to achieve low YR. Therefore, when it contains Cu, it is preferable to set it as 0.1 to 1.0% of range.
Ｍｏ：０．０１〜０．５％
Ｍｏは０．０１％以上の含有によって焼入性向上元素として作用し、多量のＭｎ添加の代替とすることができる。しかし、高価な元素であり、かつ０．５％を超えて含有しても強度上昇は飽和するため、Ｍｏを添加する場合は、０．０１〜０．５％の範囲とすることが好ましい。
Mo: 0.01 to 0.5%
When Mo is contained in an amount of 0.01% or more, it acts as a hardenability improving element and can be used as a substitute for adding a large amount of Mn. However, since it is an expensive element and the increase in strength is saturated even if it is contained in an amount exceeding 0.5%, when Mo is added, the content is preferably in the range of 0.01 to 0.5%.
Ｖ：０．０１〜０．０８％
ＶはＮｂとの複合添加により、多重溶接熱サイクル時に析出硬化するので、このような多層溶接時のＨＡＺ軟化防止に寄与する。０．０１％以上含有することで、軟化防止効果が発現するが、０．０８％を超えて含有すると析出硬化が著しくＨＡＺ靱性を劣化させるため、Ｖを含有する場合は、０．０１〜０．０８％の範囲とすることが好ましい。
V: 0.01 to 0.08%
V is precipitated and hardened during the multiple welding heat cycle due to the combined addition with Nb, which contributes to the prevention of HAZ softening during multi-layer welding. When contained in an amount of 0.01% or more, the effect of preventing softening is manifested. However, if the content exceeds 0.08%, precipitation hardening remarkably deteriorates the HAZ toughness. A range of 0.08% is preferable.
Ｂ：０．０００５〜０．００３０％
Ｂはオーステナイト粒界に偏析し、フェライト変態を抑制することで、粒界フェライト生成に起因した母材シャルピー吸収エネルギーの低下を抑制できる。この効果は、０．０００５％以上の含有により発揮されるが、０．００３０％を超えて含有してもその効果は飽和するため、Ｂを添加する場合は、０．０００５〜０．００３０％の範囲とすることが好ましい。
B: 0.0005 to 0.0030%
B segregates at austenite grain boundaries and suppresses ferrite transformation, thereby suppressing a decrease in base metal Charpy absorption energy due to the formation of grain boundary ferrite. This effect is exhibited by the content of 0.0005% or more, but even if it exceeds 0.0030%, the effect is saturated, so when adding B, 0.0005 to 0.0030% It is preferable to set it as the range.
さらに、Ｂは鋼中の固溶Ｎと結合してＢＮを形成しやすいため、十分な固溶Ｂ量を確保するために、固溶Ｎ量に対応する量を上回るようにしなくてはならない。一方、鋼中のＮは上述の通りＴｉと結合してＴｉＮを形成するため、Ｎ量と添加Ｔｉ量の化学量論比より固溶Ｎ量が計算される。よって、Ｂを含有する場合には、上記式（２）に示すようにＢ量が固溶Ｎ量（＝［Ｎ］−［Ｔｉ］／３．４）以上となるようＢ、Ｎ、Ｔｉ量を制御することが好ましい。なお、上記式（２）中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 Furthermore, since B is likely to form BN by combining with solid solution N in steel, in order to ensure a sufficient amount of solid solution B, the amount must exceed the amount corresponding to the amount of solid solution N. On the other hand, since N in steel combines with Ti to form TiN as described above, the amount of solute N is calculated from the stoichiometric ratio of the amount of N and the amount of added Ti. Therefore, when B is contained, the amounts of B, N, and Ti are set so that the amount of B becomes equal to or more than the amount of dissolved N (= [N]-[Ti] /3.4) as shown in the above formula (2). Is preferably controlled. In addition, [M] in the said Formula (2) shows the quantity (mass%) of the element M contained in steel.
Ｃａ： ０．０００５〜０．００４０％
製鋼プロセスにおいて、Ｃａ含有量が０．０００５％未満の場合、脱酸反応支配でＣａＳの確保が難しく靱性改善効果が得られず、一方、Ｃａ含有量が０．００４０％を超えた場合、粗大ＣａＯが生成しやすくなり、取鍋のノズル閉塞の原因となり、生産性を阻害する。このため、Ｃａ含有量を含有する場合には０．０００５〜０．００４０％の範囲とすることが好ましい。
Ca: 0.0005 to 0.0040%
In the steelmaking process, when the Ca content is less than 0.0005%, it is difficult to secure CaS due to the deoxidation reaction, and the toughness improving effect cannot be obtained. On the other hand, when the Ca content exceeds 0.0040%, it is coarse. CaO is likely to be generated, which causes clogging of the ladle nozzle and hinders productivity. For this reason, when it contains Ca content, it is preferable to set it as 0.0005 to 0.0040% of range.
本パラメータ式（３）は、母材のシャルピー吸収エネルギーを低下させる鋼中非金属介在物であるＭｎＳを、Ｃａ添加によって無害なＣａＯ・ＣａＳに形態制御するための鋼中Ｏ、Ｓ含有量とＣａ含有量との関係を規定したものである。式（３）を満たすことにより、粗大で靱性に悪影響を及ぼすＭｎＳ系の介在物生成を抑制するとともに、過剰なＣａ添加により生成するＣａＯ・ＣａＳの粗大化を抑制することにより、シャルピー吸収エネルギーの低下を防止する。なお、上記式中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 This parameter formula (3) is based on the contents of O and S in steel for controlling the form of MnS, which is a non-metallic inclusion in steel, which reduces the Charpy absorbed energy of the base metal, to harmless CaO / CaS by adding Ca. It defines the relationship with the Ca content. By satisfying the formula (3), the formation of MnS-based inclusions that are coarse and adversely affect toughness is suppressed, and the coarsening of CaO · CaS generated by excessive Ca addition is suppressed, thereby reducing the Charpy absorbed energy. Prevent decline. In addition, [M] in the said formula shows the quantity (mass%) of the element M contained in steel.
上記パラメータ式（３）を、より具体的に、以下に説明する。 The parameter equation (3) will be described more specifically below.
Ｃａは硫化物形成能を持ち、添加されると製鋼時の溶鋼中でシャルピー吸収エネルギーを低下させるＭｎＳの生成を抑制し、代わりに比較的靱性に無害なＣａＳを形成する。ただし、Ｃａは酸化物形成元素でもあるため、まず酸化物として消費される分を見込んだ量を添加する必要がある。すなわち、粗大で靱性に悪影響を及ぼす介在物生成抑制の観点から、Ｏ：０．００３％以下、Ｓ：０．００１％以下とした上で、ＣａＯ生成分を除いた有効ＣａＯ量（以下Ｃａ＊と記す）を実験結果の回帰による下記式（ａ）のように規定し、さらに下記式（ｂ）に示すように、ＣａとＳの化学量論比１．２５で有効Ｃａ＊を割った値が鋼中Ｓ量になるようにＣａを添加した場合、鋼中Ｓが全てＣａＳの生成に費やされる。 Ca has the ability to form sulfides, and when added, suppresses the generation of MnS, which lowers the Charpy absorbed energy in the molten steel during steelmaking, and forms CaS that is relatively harmless to toughness instead. However, since Ca is also an oxide-forming element, it is necessary to first add an amount that allows for consumption as an oxide. That is, from the viewpoint of suppressing inclusion formation that is coarse and adversely affects toughness, O: 0.003% or less, S: 0.001% or less, and effective CaO amount excluding CaO generation (hereinafter referred to as Ca *) Is defined as the following formula (a) by regression of the experimental results, and further, as shown in the following formula (b), the value obtained by dividing the effective Ca * by the stoichiometric ratio of Ca and S of 1.25 When Ca is added so that the amount of S in steel becomes the amount of S, all of the S in steel is spent on the production of CaS.
なお、下記式（ａ）、（ｂ）、（ｃ）中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 In addition, [M] in following formula (a), (b), (c) shows the quantity (mass%) of the element M contained in steel.
一方、Ｃａ含有量が過剰になると、生成するＣａＯ・ＣａＳの粗大化が生じ、シャルピー吸収エネルギーが低下することも判明した。実験室的な検討結果より、このＣａ粗大化を抑制するには、下記式（ｃ）を満たすことが求められる。 On the other hand, it was also found that when the Ca content is excessive, the CaO · CaS produced is coarsened and the Charpy absorbed energy is reduced. From the results of laboratory studies, it is required to satisfy the following formula (c) in order to suppress this Ca coarsening.
以上より、上記式（ｂ）と式（ｃ）で挟まれる範囲として上述したパラメータ式（３）が規定される。 From the above, the parameter formula (3) described above is defined as a range between the formula (b) and the formula (c).
本発明の鋼材において、上記以外の成分は、Ｆｅおよび不可避的不純物である。ただし、本発明の効果を損なわない範囲内であれば、上記以外の成分を含有することができる。 In the steel material of the present invention, components other than those described above are Fe and inevitable impurities. However, components other than those described above can be contained as long as the effects of the present invention are not impaired.
［金属組織］
金属組織はベイナイトを主体とし、この母相（第１相）であるベイナイト中に、第２相として、島状マルテンサイトが面積率で５〜１５％均一分散し、旧オーステナイト粒界に存在するフェライト相の面積率が全金属組織の５％以下とする。
[Metal structure]
The metal structure is mainly composed of bainite, and in the bainite that is the parent phase (first phase), as the second phase, island-like martensite is uniformly dispersed in an area ratio of 5 to 15% and exists in the prior austenite grain boundaries. The area ratio of the ferrite phase is 5% or less of the total metal structure.
加速冷却の冷却速度不足等でフェライト主体の組織となった場合、６００ＭＰａを超える引張強度の達成が困難となる。一方、マルテンサイト主体の組織とすると、強度は十分確保できるものの、靱性が低下する。このため、金属組織はベイナイトを主体とした組織とする。ベイナイトの面積率は、８５％以上が好ましく、９０％以上がさらに好ましい。 When the structure is mainly composed of ferrite due to insufficient cooling rate of accelerated cooling, it becomes difficult to achieve a tensile strength exceeding 600 MPa. On the other hand, if the structure is mainly composed of martensite, the toughness is lowered although sufficient strength can be secured. For this reason, the metal structure is a structure mainly composed of bainite. The area ratio of bainite is preferably 85% or more, and more preferably 90% or more.
島状マルテンサイトの面積率：５〜１５％
低降伏比を達成するため、ベイナイト中に母相となるベイナイトより硬い第２相として、島状マルテンサイトを均一に分散生成させる。島状マルテンサイトの面積率が５％未満では、降伏比が十分に低くならないため、島状マルテンサイトの面積率の下限を５％とする。
Area ratio of island martensite: 5-15%
In order to achieve a low yield ratio, island-like martensite is uniformly dispersed and formed as a second phase harder than bainite as a parent phase in bainite. If the area ratio of island martensite is less than 5%, the yield ratio is not sufficiently low. Therefore, the lower limit of the area ratio of island martensite is set to 5%.
一方、面積率が１５％を超えると、母材靱性が著しく劣化するため、島状マルテンサイトの面積率は５〜１５％の範囲とする。なお、島状マルテンサイト面積率が１５％以下であっても、例えば、旧オーステナイト粒界上にのみ存在するといったような、特定の箇所に偏在してしまうと、靱性低下をもたらすため、ベイナイトの中に均一に分散していることが好ましい。 On the other hand, if the area ratio exceeds 15%, the base material toughness deteriorates remarkably, so the area ratio of the island-like martensite is set in the range of 5 to 15%. In addition, even if the island-like martensite area ratio is 15% or less, for example, if it is unevenly distributed in a specific place such as existing only on the prior austenite grain boundary, it causes a decrease in toughness. It is preferable that it is uniformly dispersed therein.
旧オーステナイト粒界から生成したフェライトの面積率：５％以下
さらに、旧オーステナイト粒界から生成したフェライト（以下、粒界フェライトとも記す）の面積率が５％を超えると、シャルピー吸収エネルギーが著しく低下する。これは、未再結晶温度域圧延等、低温域での圧延によって、オーステナイト粒が圧延方向に伸長し、オーステナイト粒界が圧延面（鋼板面）に平行に並ぶようになり、これらのオーステナイト粒界からフェライト変態が生じると、圧延面（鋼板面）に平行にフェライト粒が並び、シャルピー衝撃試験時に延性破壊する際、主亀裂の進行に先立ち、鋼板面に平行な方向でへき開破壊を生じるため、延性亀裂伝播時のエネルギーが低下するからである。逆に、これらの鋼板面に平行に並ぶ粒界フェライトがなければ、鋼板面に平行な方向のへき開破壊が抑制されて、高い吸収エネルギーを得ることができる。従って、粒界フェライトの面積率は５％以下とする。
Area ratio of ferrite generated from prior austenite grain boundaries: 5% or less In addition, if the area ratio of ferrite generated from prior austenite grain boundaries (hereinafter also referred to as grain boundary ferrite) exceeds 5%, Charpy absorbed energy is significantly reduced. To do. This is because austenite grains are elongated in the rolling direction by rolling in a low temperature range such as non-recrystallization temperature range rolling, and the austenite grain boundaries are aligned in parallel with the rolling surface (steel plate surface). When ferrite transformation occurs, ferrite grains are aligned parallel to the rolling surface (steel plate surface), and when ductile fracture occurs during the Charpy impact test, cleavage fracture occurs in a direction parallel to the steel plate surface prior to the progress of the main crack. This is because the energy at the time of ductile crack propagation decreases. On the contrary, if there is no grain boundary ferrite aligned in parallel with these steel plate surfaces, cleavage fracture in a direction parallel to the steel plate surface is suppressed, and high absorbed energy can be obtained. Therefore, the area ratio of the grain boundary ferrite is set to 5% or less.
本発明の金属組織においては、母相であるベイナイト、および第２相である島状マルテンサイト以外の組織の面積率は小さいほどよい。ただし、前述のように、粒界フェライトは面積率が５％以下であれば許容されるほか、セメンタイトや残留オーステナイトについても、合計の面積率で１０％未満であれば、強度、靭性を損なわないので許容される。 In the metal structure of the present invention, the area ratio of the structure other than the bainite that is the parent phase and the island-like martensite that is the second phase is preferably as small as possible. However, as described above, the grain boundary ferrite is allowed as long as the area ratio is 5% or less, and the cementite and the retained austenite are not impaired in strength and toughness if the total area ratio is less than 10%. So acceptable.
セパレーションインデックス（ＳＩ）：０．０５（ｍｍ−１）以下
さらに、シャルピー衝撃試験片の破面上におけるセパレーションの指標として、本発明においては、セパレーションインデックス（ＳＩ）を用いる。その定義を図２に示す。そして、式（１）で計算されるセパレーションインデックスを０．０５（ｍｍ−１）以下とする。セパレーションの発生により、シャルピー吸収エネルギーの低下が生じるが、図３に示すシャルピー吸収エネルギーとセパレーションインデックスとの相関関係より、セパレーションインデックスが０．０５（ｍｍ−１）以下であれば、試験温度：−３０℃のシャルピー衝撃試験において、目標とする２００J以上の高い吸収エネルギーを得ることができることが判る。
［製造方法］
本発明の鋼管用鋼板は、以下の製造方法により製造することができる。
Separation index (SI): 0.05 (mm −1 ) or less Further, in the present invention, the separation index (SI) is used as an index of separation on the fracture surface of the Charpy impact test piece. The definition is shown in FIG. And the separation index calculated by Formula (1) shall be 0.05 (mm <-1> ) or less. The occurrence of separation causes a decrease in Charpy absorption energy. From the correlation between the Charpy absorption energy and the separation index shown in FIG. 3, if the separation index is 0.05 (mm −1 ) or less, the test temperature is − It can be seen that a target absorbed high energy of 200 J or more can be obtained in a 30 ° C. Charpy impact test.
[Production method]
The steel plate for steel pipes of the present invention can be manufactured by the following manufacturing method.
以下、本発明において規定される鋼の温度条件は、鋼片あるいは鋼板板厚方向平均温度を指すものとする。 Hereinafter, the temperature condition of steel defined in the present invention refers to the average temperature in the steel slab or steel plate thickness direction.
鋳造方法：連続鋳造法
ラインパイプ用鋼管素材としての経済性、生産性の観点から鋼管用鋼板の素材用である鋼片の製造は連続鋳造法とすることが好ましい。なお、鋼の製鋼方法については特に限定しないが、ラインパイプ用鋼としての経済性を確保するため、転炉法による製鋼プロセスが望ましい。
Casting method: Continuous casting method From the viewpoint of economy and productivity as a steel pipe material for line pipes, it is preferable to use a continuous casting method for the production of steel slabs for steel pipe steel materials. In addition, although it does not specifically limit about the steel making method of steel, In order to ensure the economical efficiency as steel for line pipes, the steel making process by a converter method is desirable.
鋼片加熱温度：Ａｃ３変態点以上１１００℃以下
熱間圧延を行う際、鋼片をオーステナイト化するため、鋼のオーステナイト化温度であるＡｃ３変態点以上に加熱する。なお、Ａｃ３変態点は鋼の成分組成の値を用いて、例えば下記式（４）を用いて計算することができる。一方、１１００℃を超えて加熱を行うと、結晶粒粗大化が著しく、母材ＤＷＴＴ性能の劣化が生じるため鋼片加熱温度は、Ａｃ３変態点以上１１００℃以下の範囲とする。
Steel slab heating temperature: Ac 3 transformation point or higher and 1100 ° C. or lower When hot rolling is performed, the steel slab is austenitized and heated to the Ac 3 transformation point or higher, which is the austenitizing temperature of the steel. The Ac 3 transformation point can be calculated using, for example, the following formula (4) using the value of the component composition of steel. On the other hand, when heating is performed at a temperature exceeding 1100 ° C., crystal grain coarsening is remarkable and the base material DWTT performance is deteriorated. Therefore, the billet heating temperature is set in the range of the Ac 3 transformation point to 1100 ° C.
ただし、各元素記号は各元素の含有量（質量％）を示す。 However, each element symbol indicates the content (% by mass) of each element.
熱間圧延：９５０℃以下での累積圧下率４０〜６６％
本発明において、オーステナイト再結晶温度域での熱間圧延は実施してもよいし、しなくてもよい。しかし、オーステナイト未再結晶温度域における熱間圧延は必須である。すなわち、オーステナイト未再結晶温度域である９５０℃以下で、熱間圧延を行い、オーステナイト粒を伸展させ、その後の加速冷却で変態生成するベイナイトを微細化する。ただし、本発明においては、シャルピー衝撃試験時の破面上へのセパレーション発生を抑制するため、９５０℃以下での累積圧下率の上限を６６％とする。９５０℃以下での６６％を超えるような累積圧下率で熱間圧延を施した場合、セパレーションインデックスが０．０５を超え、シャルピー吸収エネルギーが著しく低下する。なお、低温で所望のＤＷＴＴ性能を達成するため、９５０℃以下での累積圧下率を４０％以上とすることが必要であり、５０％以上とすることが好ましい。
Hot rolling: Cumulative rolling reduction at 950 ° C. or below 40 to 66%
In the present invention, hot rolling in the austenite recrystallization temperature range may or may not be performed. However, hot rolling in the austenite non-recrystallization temperature region is essential. That is, hot rolling is performed at 950 ° C. or less, which is an austenite non-recrystallization temperature range, austenite grains are extended, and bainite that is transformed by subsequent accelerated cooling is refined. However, in the present invention, in order to suppress the occurrence of separation on the fracture surface during the Charpy impact test, the upper limit of the cumulative rolling reduction at 950 ° C. or lower is set to 66%. When hot rolling is performed at a cumulative reduction rate exceeding 66% at 950 ° C. or lower, the separation index exceeds 0.05, and the Charpy absorbed energy is significantly reduced. In order to achieve the desired DWTT performance at a low temperature, the cumulative rolling reduction at 950 ° C. or lower needs to be 40% or more, and preferably 50% or more.
熱間圧延：９５０℃以下の圧延における各圧延パス圧下率１０％以下
上記オーステナイト未再結晶温度域における熱間圧延において，シャルピー衝撃試験時の破面上へのセパレーションを抑制するため、累積圧下率の上限を制限する必要があるが、さらに、各圧延パスの圧下率を低減することで、さらにセパレーションの発生を抑制する上で有効である。すなわち、各圧延パス圧下率を１０％以下にすると、セパレーションの発生やこれに伴うシャルピー吸収エネルギーの低下を抑制することができるので好ましい。
Hot rolling: Rolling pass reduction ratio of 10% or less in rolling at 950 ° C. or less In order to suppress separation on the fracture surface during Charpy impact test in hot rolling in the above austenite non-recrystallization temperature range, cumulative rolling reduction ratio Although it is necessary to limit the upper limit, the reduction of the rolling reduction of each rolling pass is effective in further suppressing the occurrence of separation. That is, it is preferable to set the rolling pass reduction ratio to 10% or less because it is possible to suppress the generation of separation and the accompanying decrease in Charpy absorbed energy.
冷却開始温度：Ａｒ３変態点温度以上
引張強度６００ＭＰａ超の高強度を達成するため、金属組織をベイナイト主体の組織にする必要がある。このため、熱間圧延後加速冷却を実施する。熱間圧延をＡｒ３変態点温度以上で終了させた上で、Ａｒ３変態点温度以上の温度から加速冷却を開始する。冷却開始温度がフェライト変態開始温度であるＡｒ３変態点を下回ると、熱間圧延後、冷却開始までの空冷過程においてオーステナイト粒界から初析フェライトが生成し、母材強度が低下するため、加速冷却を開始する温度をＡｒ３変態点以上とする。なお、Ａｒ３変態点は鋼の成分組成の値を用いて式（５）により計算することができる。
Cooling start temperature: Ar 3 transformation temperature or higher In order to achieve a high strength with a tensile strength of over 600 MPa, the metal structure must be a bainite-based structure. For this reason, accelerated cooling is performed after hot rolling. In terms of the hot rolling was terminated at Ar 3 transformation point temperature or higher and starts accelerated cooling from Ar 3 transformation point temperature or more. When the cooling start temperature falls below the Ar 3 transformation point, which is the ferrite transformation start temperature, proeutectoid ferrite is generated from the austenite grain boundaries in the air cooling process after hot rolling until the start of cooling, and the strength of the base metal is reduced. The temperature at which cooling starts is set to the Ar 3 transformation point or higher. The Ar 3 transformation point can be calculated by equation (5) using the value of the steel component composition.
ただし、各元素記号は各元素の含有量（質量％）を示す。 However, each element symbol indicates the content (% by mass) of each element.
冷却速度： １０〜８０℃／ｓ
冷却速度が１０℃／ｓ未満の場合、比較的高温で変態が進行するので、十分な強度を得ることができない。一方、８０℃／ｓを超えた冷却速度の場合、後述の冷却停止温度に制御することが難しく、特に表面近傍でマルテンサイト変態が生じ、母材靱性が著しく低下する。このため、冷却速度は１０〜８０℃／ｓの範囲とする。
Cooling rate: 10-80 ° C / s
When the cooling rate is less than 10 ° C./s, the transformation proceeds at a relatively high temperature, so that sufficient strength cannot be obtained. On the other hand, in the case of a cooling rate exceeding 80 ° C./s, it is difficult to control to the cooling stop temperature described later, and martensitic transformation occurs particularly near the surface, and the base material toughness is remarkably lowered. For this reason, a cooling rate shall be the range of 10-80 degree-C / s.
冷却停止温度： ４００〜６００℃
本発明において、加速冷却の冷却停止温度の規定は、所望の金属組織（ミクロ組織）を得るために重要である。加速冷却の停止後に行う後述の再加熱処理によりＣの濃縮した未変態オーステナイトを生成させ、この未変態オーステナイトをその後の空冷時に島状マルテンサイトへと変態させるため、ベイナイト変態途中で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４００℃未満では、ベイナイト変態が完了するため、後述の再加熱後の空冷時に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。一方、冷却停止温度が６００℃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、その後に再加熱してから冷却しても島状マルテンサイトが生成しない。このため、冷却停止温度は４００〜６００℃の範囲とする。
Cooling stop temperature: 400-600 ° C
In the present invention, the definition of the cooling stop temperature for accelerated cooling is important for obtaining a desired metal structure (microstructure). In order to produce untransformed austenite enriched with C by the reheating treatment described later after stopping the accelerated cooling, and to transform this untransformed austenite into island-like martensite during the subsequent air cooling, Cooling must be stopped in the existing temperature range.
If the cooling stop temperature is less than 400 ° C., the bainite transformation is completed, so that island-like martensite is not generated during air cooling after reheating described later, and a low yield ratio cannot be achieved. On the other hand, when the cooling stop temperature exceeds 600 ° C., C is consumed in the pearlite that precipitates during cooling, and no island-like martensite is generated even when cooled after reheating. For this reason, cooling stop temperature shall be the range of 400-600 degreeC.
再加熱温度：６００〜６８０℃
加速冷却後ただちに再加熱することで、未変態オーステナイトにＣを濃縮させその後の空冷過程で島状マルテンサイトを生成させることができる。加速冷却停止後、再加熱開始までの時間が長い場合、その間の温度低下によって未変態オーステナイトが減少し、加熱後の空冷過程で生成する島状マルテンサイト量が少なくなるため、加速冷却停止から３００秒以内で再加熱を開始することが望ましい。好ましくは１００秒以内である。
Reheating temperature: 600-680 ° C
By reheating immediately after accelerated cooling, C can be concentrated in untransformed austenite and island martensite can be generated in the subsequent air cooling process. When the time until the start of reheating is long after the accelerated cooling is stopped, untransformed austenite is reduced due to the temperature drop during that time, and the amount of island martensite generated in the air cooling process after heating is reduced. It is desirable to start reheating within seconds. Preferably, it is within 100 seconds.
さらに、再加熱温度が６００℃未満では、オーステナイトへのＣ濃化が十分に起こらず、必要とする島状マルテンサイト量を確保することができない。一方、再加熱温度が６８０℃を超えると、島状マルテンサイト量が１５％の上限を超えやすく、シャルピー吸収エネルギーの低下につながるため、再加熱温度を６００℃以上、６８０℃以下とする。 Furthermore, when the reheating temperature is less than 600 ° C., C concentration to austenite does not occur sufficiently, and the required amount of island martensite cannot be ensured. On the other hand, if the reheating temperature exceeds 680 ° C., the amount of island-like martensite tends to exceed the upper limit of 15% and leads to a decrease in Charpy absorbed energy. Therefore, the reheating temperature is set to 600 ° C. or more and 680 ° C. or less.
昇温速度が２．０℃／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ＭＡの粗大化を招く場合があり、十分な低降伏比、一様伸びを得ることができない。この機構は必ずしも明確ではないが、再加熱の昇温速度を２℃／ｓ以上と大きくすることにより、Ｃ濃縮領域の粗大化を抑制し、再加熱後の冷却過程で生成するＭＡの粗大化が抑制されるものと考えられる。 When the rate of temperature increase is less than 2.0 ° C./s, it takes a long time to reach the target reheating temperature, so that the production efficiency is deteriorated, and MA may be coarsened. Uniform elongation cannot be obtained. Although this mechanism is not necessarily clear, by increasing the heating rate of reheating to 2 ° C./s or more, the coarsening of the C-enriched region is suppressed and the coarsening of MA generated in the cooling process after reheating is increased. Is considered to be suppressed.
なお、再加熱において、加熱保持時間を設定する必要はない。また、再加熱後の冷却過程においては、冷却速度によらず島状マルテンサイトが生成するため、再加熱後の冷却条件は特に規定しないが、基本的には空冷とすることが好ましい。 In reheating, it is not necessary to set the heating holding time. Further, in the cooling process after reheating, island martensite is generated regardless of the cooling rate, and thus the cooling condition after reheating is not particularly defined, but basically it is preferably air cooling.
加速冷却後の再加熱を行うための設備として、加速冷却を行うための冷却設備の下流側に加熱装置を設置することができる。加熱装置としては、鋼板の急速加熱が可能であるガス燃焼炉や誘導加熱装置を用いることが好ましい。 As equipment for performing reheating after accelerated cooling, a heating device can be installed downstream of the cooling equipment for performing accelerated cooling. As the heating device, it is preferable to use a gas combustion furnace or induction heating device capable of rapid heating of the steel sheet.
表１に示す化学組成の鋼を用い、表２に示す熱間圧延・加速冷却・再加熱条件で鋼種Ａ〜Ｈに対してＮｏ．１〜１６の鋼板を作製した。 Using steel with the chemical composition shown in Table 1, No. was compared with the steel types A to H under the hot rolling, accelerated cooling, and reheating conditions shown in Table 2. 1 to 16 steel plates were produced.
鋼種Ａ〜Ｅは発明例、鋼種Ｆ〜Ｈは比較例である。比較例ＥはＣが、比較例ＦはＭｎが、比較例ＧはＮｂがそれぞれ本発明の範囲外となっている。 Steel types A to E are invention examples, and steel types F to H are comparative examples. In Comparative Example E, C is outside the scope of the present invention, in Comparative Example F is outside Mn, and in Comparative Example G is outside Nb.
作製された鋼板のうち、Ｎｏ．１〜９は発明例を、Ｎｏ．１０〜１６は比較例を示しており、比較例Ｎｏ．１０〜１３は製造条件のいずれかが、また、比較例Ｎｏ．１４〜１６は鋼種成分がそれぞれ発明の範囲外である。 Among the produced steel plates, No. 1 to 9 are invention examples. 10 to 16 show comparative examples. 10 to 13 are any of the production conditions. As for 14-16, a steel seed component is outside the range of an invention, respectively.
表２の製造方法により得られた鋼板の板幅中央部より金属組織観察用サンプルを採取し、圧延長手方向と平行な板厚断面を鏡面研磨したあと、２段エッチング法を用いて島状マルテンサイトを出現させた。 A sample for metallographic observation was taken from the center of the plate width of the steel plate obtained by the manufacturing method shown in Table 2, and after mirror-polishing the plate thickness section parallel to the rolling longitudinal direction, the island shape was obtained using a two-step etching method. Martensite appeared.
その後、走査型電子顕微鏡（ＳＥＭ）を用い、２０００倍の倍率で無作為に１０視野について金属組織写真を撮影し、写真中の島状マルテンサイトの面積率を画像解析装置にて測定した。 Then, using a scanning electron microscope (SEM), metal structure photographs were taken for 10 visual fields at a magnification of 2000 times, and the area ratio of island martensite in the photographs was measured with an image analyzer.
さらに、同試料を再鏡面研磨し、ナイタールエッチングを行ってから、光学顕微鏡を用いて４００〜１０００倍の範囲の倍率で無作為に５視野写真撮影を行い、画像解析処理によって粒界フェライト組織の面積率を算出した。 Furthermore, after re-polished the sample and performing nital etching, 5 field photographs were randomly taken at a magnification of 400 to 1000 times using an optical microscope, and the grain boundary ferrite structure was obtained by image analysis processing. The area ratio was calculated.
次に、それぞれの鋼板より圧延方向と直交する方向が、試験片長手方向となるようにＡＰＩ−５Ｌに準拠した全厚引張試験片およびＤＷＴＴ試験片、および板厚中央位置からＪＩＳ Ｚ２２０２（１９８０改訂版）のＶノッチシャルピー衝撃試験片を採取し、鋼板の引張試験、ＤＷＴＴ試験、およびシャルピー衝撃試験（ＪＩＳ Ｚ２２４２）を実施して、強度と靱性を評価した。 Next, JIS Z2202 (1980 revision) from the full thickness tensile test piece and DWTT test piece based on API-5L and the plate thickness center position so that the direction orthogonal to the rolling direction is the test piece longitudinal direction from each steel plate. Plate) V-notch Charpy impact test pieces were collected and subjected to tensile test, DWTT test, and Charpy impact test (JIS Z2242) of steel plates to evaluate strength and toughness.
次に、シャルピー衝撃試験後の残片より破面部の外観写真を１０倍の倍率で撮影し、その画像から、破面中に存在する、１ｍｍ以上の長さを持つ個々のセパレーション長さを計測し、これを元にしてセパレーションインデックスを求めた。 Next, a photograph of the appearance of the fractured surface portion was taken from the remaining pieces after the Charpy impact test at a magnification of 10 times, and the length of each separation having a length of 1 mm or more existing in the fractured surface was measured from the image. Based on this, a separation index was obtained.
母材の金属組織の画像解析結果および強度・靱性調査結果をまとめて表３に示す。 Table 3 summarizes the results of the image analysis of the metal structure of the base metal and the results of the strength / toughness investigation.
本発明例であるＮｏ．１〜９は、いずれも本発明の鋼板化学組成、ミクロ組織（金属組織）、圧延・加速冷却・再加熱条件範囲内であり、目標とする引張強度６００ＭＰａ以上、降伏比８０％以下、−２０℃でのＤＷＴＴ延性破面率８５％以上、および−３０℃でのシャルピー吸収エネルギー２００Ｊ以上を満足した。 No. which is an example of the present invention. 1 to 9 are all in the chemical composition, microstructure (metal structure), rolling / accelerated cooling / reheating condition range of the present invention, and target tensile strength is 600 MPa or more, yield ratio is 80% or less, −20 The DWTT ductile fracture surface ratio at 85 ° C. was 85% or more, and Charpy absorbed energy at −30 ° C. was 200 J or more.
一方、スラブ加熱温度が、本発明の範囲の上限を上回った比較例であるＮｏ．１０は、ＤＷＴＴ性能が目標を下回った。９５０℃以下の熱間圧延での累積圧下率が上限を超えた比較例であるＮｏ．１１は、セパレーションの発生が著しく、シャルピー吸収エネルギーが目標を下回った。逆に、９５０℃以下の熱間圧延での累積圧下率が下限を下回った比較例であるＮｏ．１２は、ＤＷＴＴ性能が目標を下回った。加速冷却の冷却開始温度が下限を下回った比較例であるＮｏ．１３は、旧オーステナイト粒界上に多量のフェライトが生成した結果、セパレーションの発生が著しく、シャルピー吸収エネルギーが目標を下回った。 On the other hand, No. 1 which is a comparative example in which the slab heating temperature exceeded the upper limit of the range of the present invention. 10 DWTT performance was below the target. A comparative example in which the cumulative rolling reduction in hot rolling at 950 ° C. or lower exceeded the upper limit. In No. 11, separation occurred remarkably and Charpy absorbed energy was lower than the target. On the other hand, No. which is a comparative example in which the cumulative rolling reduction in hot rolling at 950 ° C. or lower is below the lower limit. 12, DWTT performance was below target. A comparative example in which the cooling start temperature of accelerated cooling was below the lower limit. In No. 13, a large amount of ferrite was generated on the prior austenite grain boundaries, and as a result, separation occurred remarkably and the Charpy absorbed energy was below the target.
また、母材成分のＣ量が本発明の下限を下回った比較例であるＮｏ．１４は、金属組織中の島状マルテンサイト分率が必要量確保できなかったため、低降伏比が得られなかった。逆に、母材成分のＭｎ量、Ｎｂ量が本発明の範囲の上限を上回った比較例であるＮｏ．１５およびＮｏ．１６は、島状マルテンサイト分率が上限を超えたため、母材シャルピー吸収エネルギーが目標を下回った。 In addition, in the comparative example in which the amount of C of the base material component was below the lower limit of the present invention, No. No. 14 could not secure a low yield ratio because the required amount of island martensite fraction in the metal structure could not be secured. On the contrary, No. 1 which is a comparative example in which the amount of Mn and the amount of Nb of the base material component exceeded the upper limit of the range of the present invention. 15 and no. For No. 16, the island-shaped martensite fraction exceeded the upper limit, so the base material Charpy absorbed energy was below the target.
Claims (6)
なお、上記式中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 Furthermore, it contains B: 0.0005-0.0030% by mass%, and B, Ti, and N satisfy | fill following formula (2), The high of Claim 1 or 2 characterized by the above-mentioned. Steel sheet for tough and highly deformable high strength steel pipes.
In addition, [M] in the said formula shows the quantity (mass%) of the element M contained in steel.
なお、上記式中の［Ｍ］は、鋼中に含有される元素Ｍの量（質量％）を示す。 Furthermore, it contains Ca: 0.0005-0.0040% by mass%, and Ca, O, and S satisfy | fill following formula (3), Any one of Claim 1 thru | or 3 characterized by the above-mentioned. Steel sheet for high toughness and high deformability high strength steel pipe as described.
In addition, [M] in the said formula shows the quantity (mass%) of the element M contained in steel.
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