JP2008169467A

JP2008169467A - High-strength thick steel plate having excellent brittle crack propagation-stopping performance, and method for producing the same

Info

Publication number: JP2008169467A
Application number: JP2007228112A
Authority: JP
Inventors: Kiyotaka Nakajima; 清孝中島; Masanori Minagawa; 昌紀皆川; Akira Shishibori; 明獅々堀
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 2006-12-14
Filing date: 2007-09-03
Publication date: 2008-07-24
Anticipated expiration: 2027-09-03
Also published as: JP5064149B2

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-strength thick steel plate having excellent brittle crack propagation-stopping performance, and to provide a method for producing the same. <P>SOLUTION: The high-strength thick steel plate contains as chemical components, by mass, 0.04 to 0.15% C, 0.1 to 0.5% Si, 0.5 to 2.5% Mn, ≤0.02% P, ≤0.01% S, 0.001 to 0.1% Al, 0.005 to 0.02% Ti and 0.001 to 0.008% N, and the balance iron with inevitable impurities, and the microstructure is a ferrite or/and pearlite structure where bainite is used as a host phase. In the case the structure is divided into the three layers of front-rear layer parts having 25% in depth from the front-rear faces of the steel sheet and a sheet thickness central part other than that, the front-rear layer parts has a texture where the (100) X-ray face intensity ratio parallel to the rolling face is 1.5 to <2.0 is provided in the region of 5 to 25% of the sheet thickness in the surface-back layer parts, and the sheet thickness central part has a texture where the (111) or/and (211) X-ray face intensity ratio parallel to the rolling face is ≥2.0. <P>COPYRIGHT: (C)2008,JPO&INPIT

Description

本発明は、脆性き裂伝播停止性能に優れた高強度厚鋼板及びその製造方法に関する。 The present invention relates to a high-strength thick steel plate excellent in brittle crack propagation stopping performance and a method for producing the same.

造船、建築、タンク、海洋構造物、ラインパイプなどの溶接構造物に用いられる厚鋼板には、構造物の致命的な破壊を防止するために、脆性破壊が伝播することを停止する能力である脆性き裂伝播停止性能（アレスト性）が求められる。近年、構造物の大型化に伴い、降伏応力３５５〜５５０ＭＰａ、板厚４０〜１００ｍｍの高強度厚鋼板を使用するケースが多くなっている。しかし、上記したアレスト性は、一般に強度及び板厚それぞれに相反する傾向にある。このため、高強度厚鋼板においてアレスト性を向上させる技術が望まれている。 Thick steel plates used in shipbuilding, construction, tanks, offshore structures, line pipes and other welded structures have the ability to stop the propagation of brittle fracture to prevent catastrophic destruction of the structure. Brittle crack propagation stopping performance (arrestability) is required. In recent years, with the increase in size of structures, there are increasing cases of using high-strength thick steel plates having a yield stress of 355 to 550 MPa and a plate thickness of 40 to 100 mm. However, the above-described arrestability generally tends to conflict with strength and thickness. For this reason, the technique which improves arrestability in a high intensity | strength thick steel plate is desired.

アレスト性を向上させる方法として、例えば結晶粒径を制御する方法、脆化第二相を制御する方法、及び集合組織を制御する方法が知られている。 As a method for improving the arrestability, for example, a method of controlling the crystal grain size, a method of controlling the embrittled second phase, and a method of controlling the texture are known.

結晶粒径を制御する方法としては、特許文献１に記載された技術、及び特許文献２、３に記載された技術がある。特許文献１に記載された技術は、フェライトを母相としたものであり、このフェライトを細粒化することにより、アレスト性を向上させるものである。そのような細粒フェライトを得るために、表裏層部より鋳片厚中心方向に鋳片厚の１／８以上がＡｒ３以下となるように冷却し、極低温域で圧延を行い、その後Ａｃ３を越える温度まで復熱させ、フェライトを再結晶させる必要がある。特許文献２、３に記載された技術は、フェライトを母相としたものであり、表層部を一旦Ａｒ１以下に冷却し、その後表層部が復熱する過程で圧延を行うことにより、微細なフェライト再結晶粒を得るものである。 As a method for controlling the crystal grain size, there are a technique described in Patent Document 1 and a technique described in Patent Documents 2 and 3. The technique described in Patent Document 1 uses ferrite as a parent phase, and improves the arrestability by making the ferrite finer. In order to obtain such fine-grained ferrite, cooling is performed so that 1/8 or more of the slab thickness becomes Ar3 or less from the front and back layer portions toward the center of the slab thickness, and rolling is performed in a cryogenic region, and then Ac3 is used. It is necessary to reheat to a temperature exceeding that and recrystallize the ferrite. The technologies described in Patent Documents 2 and 3 use ferrite as a parent phase. The surface layer portion is once cooled to Ar1 or less, and then rolled in the process of reheating the surface layer portion, whereby fine ferrite is obtained. Recrystallized grains are obtained.

また、脆化第二相を制御する方法としては、特許文献４に記載された技術がある。特許文献４に記載された技術は、母相となるフェライト中に微細な脆化第二相（例えばマルテンサイト）を分散させることにより、脆性き裂先端部において脆化第二相に微小き裂を発生させて、き裂先端の応力状態を緩和させるものである。 As a method for controlling the embrittled second phase, there is a technique described in Patent Document 4. In the technique described in Patent Document 4, a fine embrittled second phase (for example, martensite) is dispersed in ferrite as a parent phase, whereby a microcrack is formed in the embrittled second phase at the brittle crack tip. To reduce the stress state at the crack tip.

さらに、集合組織を制御する方法としては、特許文献５に記載された技術がある。特許文献５に記載された技術は、極低炭素（Ｃ＜０．０３％）のベイナイト単相鋼において、圧延面と平行に（２１１）集合組織を発達させるものである。 Furthermore, as a method for controlling the texture, there is a technique described in Patent Document 5. The technique described in Patent Document 5 is to develop a (211) texture parallel to the rolling surface in a bainite single-phase steel with extremely low carbon (C <0.03%).

特開昭６１−２３５５３４号公報JP 61-235534 A 特開２００３−２２１６１９号公報JP 2003-221619 A 特開平５−１４８５４２号公報JP-A-5-148542 特開昭５９−４７３２３号公報JP 59-47323 A 特開２００２−２４１８９１号公報JP 2002-241891 A

特許文献１〜３に記載の技術では、フェライトの再結晶を利用してフェライトを母相にしているため、高強度で、かつ、板厚の厚い鋼板とすることが困難である。また、冷却、圧延、復熱工程を経る必要があり、製造プロセスが複雑になるため、安定した材質を得るのは極めて困難である。さらに、このような製造プロセスでは、板面の冷却が不均一になることに起因した形状不良が生じやすい。形状不良が生じた場合、形状矯正に多大なコストを要する。 In the techniques described in Patent Documents 1 to 3, since ferrite is used as a parent phase by utilizing recrystallization of ferrite, it is difficult to obtain a steel plate having high strength and a large plate thickness. In addition, it is necessary to go through cooling, rolling, and recuperation steps, and the manufacturing process becomes complicated. Therefore, it is extremely difficult to obtain a stable material. Furthermore, in such a manufacturing process, shape defects are likely to occur due to uneven cooling of the plate surface. When a shape defect occurs, a great deal of cost is required for shape correction.

また、特許文献４に記載の技術では、フェライト中にマルテンサイト分散させているので脆性き裂発性特性が著しく劣化してしまう。さらに、フェライトを母相としているため、上記同様に高強度かつ板厚が厚い鋼板とすることが困難である。 Further, in the technique described in Patent Document 4, since martensite is dispersed in ferrite, the brittle cracking property is remarkably deteriorated. Furthermore, since ferrite is used as a parent phase, it is difficult to obtain a steel plate having high strength and a thick plate thickness as described above.

また、特許文献５に記載の技術では、極低炭素ベイナイト単相鋼にする必要がある。このような極低炭素鋼を得るためには、転炉内での酸素吹き付けのみでは脱炭が不十分であり、真空脱ガス工程での脱炭工程が追加されることなり、製鋼コストが増加する。また融点が上昇するため、溶鋼温度を上げる必要があることから、耐火物が劣化する原因となり、製鋼負荷が極めて大きい。さらに極低炭素鋼では強度を確保するために高合金にする必要があるので、合金コストが増加する。また高合金であるため溶接熱影響部（ＨＡＺ）の靭性が劣化してしまう。そして、ベイナイト単相組織形成によって板厚方向に均一な集合組織を発達させる方法では、アレスト性を飛躍的に向上させることはできない。 Moreover, in the technique described in Patent Document 5, it is necessary to use an extremely low carbon bainite single phase steel. In order to obtain such an ultra-low carbon steel, decarburization is not sufficient only by oxygen blowing in the converter, and a decarburization process in the vacuum degassing process is added, which increases the steelmaking cost. To do. Moreover, since melting | fusing point rises and it is necessary to raise molten steel temperature, it becomes a cause of refractory deterioration and the steelmaking load is very large. Furthermore, since it is necessary to make it a high alloy in order to ensure intensity | strength in ultra-low carbon steel, an alloy cost increases. Moreover, since it is a high alloy, the toughness of a welding heat affected zone (HAZ) will deteriorate. And, by the method of developing a uniform texture in the thickness direction by forming a bainite single phase structure, the arrestability cannot be improved dramatically.

本発明は上記のような事情を考慮してなされたものであり、その課題は、製造コストが低く、強度が高く、ＨＡＺ靭性の劣化がない、脆性き裂伝播停止性能に優れた高強度厚鋼板及びその製造方法を提供することにある。 The present invention has been made in consideration of the above-mentioned circumstances, and the problem is that the manufacturing cost is low, the strength is high, the HAZ toughness is not deteriorated, and the brittle crack propagation stopping performance is excellent. It is in providing a steel plate and its manufacturing method.

高強度厚鋼板のアレスト性を向上させるためには板厚方向に集合組織が異なるような分布を形成させることによって、き裂伝播の抵抗となるような破面を形成させる集合組織制御が必要である。このような集合組織制御技術に関し、本発明者らが鋭意検討した結果、鋼板の表面及び裏面から板厚の２５％までの表裏層部とそれ以外の板厚中心部の三層に分けたとき、表裏層部の領域において、板厚の５％以上２５％以下の厚さで圧延面と平行な（１００）Ｘ線面強度比が１．５以上２．０未満の集合組織が形成し、かつ、板厚中心部において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比が２．０以上の集合組織を形成させると、表裏層部の領域が脆性き裂の伝播抵抗となりアレスト性が飛躍的に向上すること見出して、本発明を完成した。本発明の要旨は、以下の通りである。 In order to improve the arrestability of high-strength thick steel plates, it is necessary to control the texture to form a fracture surface that can resist crack propagation by forming a distribution with a different texture in the thickness direction. is there. As a result of intensive studies by the present inventors regarding such a texture control technique, when the front and back layer portions from the front and back surfaces of the steel plate to 25% of the plate thickness and the other three thickness center portions are divided. In the region of the front and back layers, a texture with a thickness of 5% or more and 25% or less of the plate thickness parallel to the rolling surface and having a (100) X-ray plane intensity ratio of 1.5 or more and less than 2.0 is formed, In addition, when a texture having a (111) or / and (211) X-ray plane intensity ratio of 2.0 or more is formed in the center portion of the plate thickness, the region of the front and back layer portions propagates a brittle crack. The present invention was completed by finding that the resistance and the arrestability were dramatically improved. The gist of the present invention is as follows.

（１）質量％で、
Ｃ：０．０４〜０．１５％、
Ｓｉ：０．１〜０．５％、
Ｍｎ：０．５〜２．５％、
Ｐ：≦０．０２％、
Ｓ：≦０．０１％、
Ａｌ：０．００１〜０．１％、
Ｔｉ：０．００５〜０．０２％、
Ｎ：０．００１〜０．００８％
を含有し、残部が鉄及び不可避不純物によって化学成分が構成された鋼板で、ミクロ組織がベイナイトを母相としたフェライト又は／及びパーライト組織であり、鋼板の表面及び裏面から板厚の２５％までの表裏層部とそれ以外の板厚中心部の三層に分けたとき、該表裏層部で、板厚の５％以上２５％以下の厚さに圧延面と平行な（１００）Ｘ線面強度比が１．５以上２．０未満の集合組織を有しており、それ以外の板厚中心部の領域において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比が２．０以上の集合組織を有していることを特徴とする脆性き裂伝播停止性能に優れた高強度厚鋼板。 (1) In mass%,
C: 0.04 to 0.15%,
Si: 0.1 to 0.5%,
Mn: 0.5 to 2.5%
P: ≦ 0.02%,
S: ≦ 0.01%
Al: 0.001 to 0.1%,
Ti: 0.005 to 0.02%,
N: 0.001 to 0.008%
The balance is a steel plate in which the chemical composition is composed of iron and inevitable impurities, and the microstructure is a ferrite or / and pearlite structure with bainite as a parent phase, from the front and back surfaces of the steel plate to 25% of the plate thickness (100) X-ray plane parallel to the rolling surface at a thickness of 5% to 25% of the plate thickness at the front and back layer portions. (111) or / and (211) X-ray plane intensity ratio parallel to the rolling surface has a texture with an intensity ratio of 1.5 or more and less than 2.0, and in the other region of the center of the plate thickness. A high-strength thick steel plate excellent in brittle crack propagation stopping performance, characterized by having a texture of 2.0 or more.

（２）質量％で、
Ｃｕ：０．０５〜１％、
Ｎｉ：０．０５〜２％、
Ｃｒ：０．０５〜１％、
Ｍｏ：０．０５〜０．５％、
Ｎｂ：０．００３〜０．１％、
Ｖ：０．００５〜０．２％、
Ｂ：０．０００２〜０．００３％
の少なくとも１種以上を化学成分として含有することを特徴とする上記（１）に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 (2) In mass%,
Cu: 0.05 to 1%,
Ni: 0.05-2%,
Cr: 0.05 to 1%,
Mo: 0.05-0.5%
Nb: 0.003 to 0.1%,
V: 0.005 to 0.2%,
B: 0.0002 to 0.003%
The high-strength thick steel plate having excellent brittle crack propagation stopping performance as described in (1) above, comprising at least one of the above as chemical components.

（３）質量％で、
Ｃａ：０．０００３〜０．００５％、
Ｍｇ：０．０００３〜０．００５％、
ＲＥＭ：０．０００３〜０．００５％
の少なくとも１種以上を化学成分として含有することを特徴とする上記（１）または（２）に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 (3) In mass%,
Ca: 0.0003 to 0.005%,
Mg: 0.0003 to 0.005%,
REM: 0.0003 to 0.005%
A high-strength thick steel plate having excellent brittle crack propagation stopping performance as described in (1) or (2) above, which contains at least one of the above as chemical components.

（４）板厚が４０ｍｍ以上であることを特徴とする上記（１）〜（３）のいずれか１項に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 (4) The high-strength thick steel plate excellent in brittle crack propagation stopping performance according to any one of (1) to (3) above, wherein the plate thickness is 40 mm or more.

（５）前記表裏層部が分岐き裂となりに、該分岐き裂が板厚の５％以上２５％以下の長さで外部応力と垂直方向に伝播し、それ以外の板厚中心部の領域が主き裂となり外部応力と垂直な面に対し１５°以上４５°以下で伝播することを特徴とする上記（１）〜（４）のいずれかに記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 (5) When the front and back layer portions become branched cracks, the branched cracks propagate in the direction perpendicular to the external stress with a length of 5% or more and 25% or less of the plate thickness. It is excellent in brittle crack propagation stopping performance according to any one of the above (1) to (4), characterized in that it becomes a main crack and propagates at 15 ° to 45 ° with respect to a plane perpendicular to external stress. High strength thick steel plate.

（６）上記（１）〜（３）のいずれかに記載の化学成分を有する鋼片を、９５０〜１２５０℃に加熱し、表面温度を６５０℃以上８５０℃以下、板厚中心温度を８５０℃超１０５０℃以下で累積圧下率３０％以上の圧延を行った後、表面温度６００℃以上から、板厚平均で８℃/s以上の冷却速度で５００℃以下の温度まで加速冷却を行うことを特徴とする脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 (6) The steel slab having the chemical component according to any one of (1) to (3) is heated to 950 to 1250 ° C, the surface temperature is 650 ° C to 850 ° C, and the plate thickness center temperature is 850 ° C. After rolling at an ultra-low temperature of 1050 ° C. or less and a cumulative reduction ratio of 30% or more, accelerated cooling is performed from a surface temperature of 600 ° C. or more to a temperature of 500 ° C. or less at a cooling rate of 8 ° C./s or more on the average thickness. A method for producing a high strength thick steel plate having excellent brittle crack propagation stopping performance.

（７）前記加速冷却を終了した後、３００℃以上６５０℃以下で焼戻しすることを特徴とする上記（６）に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 (7) The method for producing a high-strength thick steel plate excellent in brittle crack propagation stopping performance as described in (6) above, wherein after accelerating cooling is completed, tempering is performed at 300 ° C. or more and 650 ° C. or less.

（８）板厚が４０ｍｍ以上であることを特徴とする上記（６）または（７）に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 (8) The method for producing a high-strength thick steel plate having excellent brittle crack propagation stopping performance as described in (6) or (7) above, wherein the plate thickness is 40 mm or more.

本発明鋼によれば、脆性き裂伝播停止性能に極めて優れ、かつ、強度が高く、板厚が大きく、ＨＡＺ靭性の劣化がない鋼板となるので、溶接鋼構造物の低コスト化や安全性向上を図れる。さらに、本発明の製造方法では、製鋼、圧延、精製工程の負荷が少なく、低コストで、生産性良く、安定して製造可能となる。 According to the steel of the present invention, since the steel sheet is extremely excellent in brittle crack propagation stopping performance, has high strength, has a large plate thickness, and does not deteriorate the HAZ toughness, it is possible to reduce the cost and safety of the welded steel structure. Improvements can be made. Furthermore, in the production method of the present invention, the steelmaking, rolling, and refining processes are less burdened, and the production can be stably performed at low cost with high productivity.

以下、本発明の実施形態について説明する。本実施形態に係る高強度厚鋼板は、ミクロ組織がベイナイトを母相としたフェライト又は／及びパーライト組織であり、かつ板厚方向の集合組織分布を制御することにより、脆性き裂伝播停止性能を向上させるものである。 Hereinafter, embodiments of the present invention will be described. The high-strength thick steel plate according to this embodiment is a ferrite or / and pearlite structure whose microstructure is bainite as a parent phase, and has a brittle crack propagation stopping performance by controlling the texture distribution in the thickness direction. It is to improve.

鋼板の脆性き裂は圧延面と平行な（１００）のへき開面に沿って伝播することから、特定の集合組織を発達させることによってき裂の伝播方向を変化させることが可能である。ただし、板厚方向に均一な集合組織が形成されたとしてもマクロなき裂伝播方向を変化させることはできるが、アレスト性を飛躍的に向上させるには至らない。このアレスト性を向上させるためには板厚方向に集合組織が異なるような分布を形成させることによって、き裂伝播の抵抗となるような破面を形成させることが効果的である。このような集合組織制御技術に関し、本発明者らが鋭意検討した結果、鋼板の表面及び裏面から板厚の２５％までの表裏層部とそれ以外の板厚中心部の三層に分けたとき、該表裏層部で、板厚の５％以上２５％以下の厚さに圧延面と平行な（１００）Ｘ線面強度比が１．５以上２．０未満の集合組織形成させ、かつ、前記板厚中心部において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比が２．０以上の集合組織を形成させる、即ち、圧延面と平行な（１００）Ｘ線面強度比が１．５以上２．０未満の集合組織を、表裏層部で、少なくとも鋼板の表裏面から板厚の５％またはこれを超え、多くとも板厚の２５％の厚さで形成させ、更に、圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比が２．０以上の集合組織を、この表裏層部より鋼板中心側で、板厚中心（ｔ／２）を含む板厚中心部領域に形成させることによって、脆性き裂伝播を増大させアレスト性が飛躍的に向上することが明らかとなった。 Since a brittle crack in a steel plate propagates along a (100) cleavage plane parallel to the rolling surface, it is possible to change the propagation direction of the crack by developing a specific texture. However, even if a uniform texture is formed in the plate thickness direction, the macro crack propagation direction can be changed, but the arrestability cannot be improved dramatically. In order to improve the arrestability, it is effective to form a fracture surface that provides resistance to crack propagation by forming a distribution with different textures in the thickness direction. As a result of intensive studies by the present inventors regarding such a texture control technique, when the front and back layer portions from the front and back surfaces of the steel plate to 25% of the plate thickness and the other three thickness center portions are divided. In the front and back layer portions, a (100) X-ray plane intensity ratio parallel to the rolling surface is formed to a thickness of 5% or more and 25% or less of the plate thickness, and a texture of 1.5 or more and less than 2.0 is formed, and The (111) or / and (211) X-ray surface intensity ratio parallel to the rolling surface is formed at the center of the plate thickness to form a texture with 2.0 or more, that is, the (100) X-ray surface parallel to the rolling surface. A texture having a strength ratio of 1.5 or more and less than 2.0 is formed at the front and back layer portions at least 5% or more of the plate thickness from the front and back surfaces of the steel plate and at most 25% of the plate thickness. Furthermore, the (111) or / and (211) X-ray surface intensity ratio parallel to the rolling surface has a texture of 2.0 or more. It is clear that by forming in the plate thickness center region including the plate thickness center (t / 2) on the steel plate center side from the front and back layer portions, brittle crack propagation is increased and arrestability is dramatically improved. became.

本発明者は、このアレスト性が飛躍的に向上する原因を解明するため、先ず、アレスト性の優れた鋼板（−１０℃のときのアレスト靭性値Ｋｃａが６５００Ｎ/ｍｍ^1.5で板厚６０ｍｍの鋼板）を用いて、温度勾配型の標準ＥＳＳＯ試験を行った。
尚、温度勾配型の標準ＥＳＳＯ試験とは「ＷＥＳ鋼種認定試験方法」（２０００年３月改正版）の「７．脆性破壊伝播停止試験」に準拠する試験である。
この試験結果より、図１（ａ）に示す様に、予め形成した切欠き（予切欠き）２を有する鋼板Ｔが外部からの引張応力(以下単に外部応力と称す)に対し、この予切欠き２から発生したマクロな脆性き裂３は、外部応力１に対して３０°の傾斜角度θを持って矢印4方向に伝播し、そして、表層部分には外部応力１と直角方向（矢印６方向）に伝播している前記主き裂３から分岐した長さ数ｍｍの小さなき裂（以下、単に分岐き裂と称す）５が多数観察された。 In order to elucidate the cause of the dramatic improvement in arrestability, the present inventor firstly made a steel plate having excellent arrestability (a steel plate having an arrest toughness value Kca at −10 ° C. of 6500 N / mm ^1.5 and a thickness of 60 mm). ) Was used to perform a temperature gradient type standard ESSO test.
The temperature gradient type standard ESSO test is a test that conforms to “7. Brittle fracture propagation stop test” of “WES steel type certification test method” (revised in March 2000).
From this test result, as shown in FIG. 1 (a), the steel plate T having a notch (pre-notch) 2 formed in advance is subjected to this pre-cut against the external tensile stress (hereinafter simply referred to as external stress). The macro brittle crack 3 generated from the notch 2 propagates in the direction of arrow 4 with an inclination angle θ of 30 ° with respect to the external stress 1, and in the direction perpendicular to the external stress 1 (arrow 6). Many small cracks (hereinafter simply referred to as “branch cracks”) 5 having a length of several millimeters branched from the main crack 3 propagating in the direction) were observed.

また、板厚方向（図１のＡ−Ａ矢視方向）からこの破面を観察すると図１（ｂ）に示す様に、このマクロな脆性き裂３は、板厚方向中心部９に発生する主き裂８と、該主き裂８から分岐して鋼板表裏面まで達する前記分岐き裂５から構成されていることが観察された。そして、この主き裂８は中心部の３６ｍｍの厚さ（板厚の６０％）であり、分岐き裂５の長さは鋼板表裏面で各々１２ｍｍ（板厚の２０％）であった。また、この分岐き裂５は図１（ａ）で、マクロき裂の発生位置より上方側に多く発生し、下方側には殆ど発生しないことが判った。 Further, when this fracture surface is observed from the thickness direction (the direction of arrows AA in FIG. 1), as shown in FIG. 1B, this macro brittle crack 3 is generated at the central portion 9 in the thickness direction. It was observed that the main crack 8 and the branched crack 5 branched from the main crack 8 and reached the front and back surfaces of the steel sheet were observed. The main crack 8 had a thickness of 36 mm at the center (60% of the plate thickness), and the lengths of the branched cracks 5 were 12 mm (20% of the plate thickness) on the front and back surfaces of the steel plate. Further, in FIG. 1 (a), it was found that a lot of the branch cracks 5 are generated above the macro crack generation position and hardly occur on the lower side.

次に、複数のアレスト靭性値Ｋｃａの優れた鋼板（Ｋｃａ：―１０℃で５０００Ｎ／ｍｍ^1.5以上の鋼板）とアレスト靭性値Ｋｃａの劣る鋼板（Ｋｃａ：―１０℃で５０００Ｎ／ｍｍ^1.5未満の鋼板）について、前記主き裂８の傾斜角度θと分岐き裂５の発生位置との関係について調査するために、前記同様の温度勾配型の標準ＥＳＳＯ試験を行った。この結果を図２に示す。尚、試験に使用した鋼板は板厚４０〜１００ｍｍ、降伏応力３６０〜５４０ＭＰａであった。 Next, a plurality of steel plates having excellent arrest toughness value Kca (Kca: steel plate having 5000 N / mm ^1.5 or more at −10 ° C.) and steel plates having inferior arrest toughness value Kca (Kca: steel plate having less than 5000 N / mm ^1.5 at −10 ° C. In order to investigate the relationship between the inclination angle θ of the main crack 8 and the generation position of the branch crack 5, the same temperature gradient type standard ESSO test as described above was performed. The result is shown in FIG. In addition, the steel plate used for the test had a plate thickness of 40 to 100 mm and a yield stress of 360 to 540 MPa.

この図２から、前記アレスト性の優れた鋼板は、主き裂８の傾斜角度θは１５°以上、４５°以下の範囲で、分岐き裂は板厚の５以上、２５％以下の長さであることが判る。一方、アレスト靭性値Ｋｃａの劣る鋼板は、分岐き裂が発生せずに主き裂のみである場合、または、主き裂８の傾斜角度θは１５°以上の場合であるが、分岐き裂が板厚の５％未満の長さである事が判る。また、主き裂８の傾斜角度θが１５°未満になる場合には分岐き裂５が発生する鋼板はなかった。また、分岐き裂５が板厚の２５％超の長さに及んで発生している鋼板もなかった。
このことから、アレスト性の優れた鋼板、即ち、アレスト靭性値Ｋｃａが―１０℃で５０００Ｎ／ｍｍ^1.5以上を示す鋼板は、主き裂８の傾斜角度θが１５°以上、４５°以下で、かつ、分岐き裂が板厚の５％以上２５％以下の長さで鋼板の表裏層部７に発生することが判る。 As shown in FIG. 2, the steel plate having excellent arrestability has an inclination angle θ of the main crack 8 in the range of 15 ° to 45 °, and the branch crack has a length of 5 to 25% of the plate thickness. It turns out that it is. On the other hand, a steel plate having an inferior arrest toughness value Kca is a case where only a main crack is generated without a branch crack, or a case where the inclination angle θ of the main crack 8 is 15 ° or more. Is less than 5% of the plate thickness. Further, when the inclination angle θ of the main crack 8 was less than 15 °, there was no steel plate in which the branch crack 5 occurred. Further, there was no steel plate in which the branch crack 5 occurred over a length exceeding 25% of the plate thickness.
From this, a steel plate having excellent arrestability, that is, a steel plate having an arrest toughness value Kca of 5,000 ° C. at -10 ° C. of ^1.5 N or more is an inclination angle θ of the main crack 8 of 15 ° or more and 45 ° or less, And it turns out that a branch crack generate | occur | produces in the front-and-back layer part 7 of a steel plate with the length of 5 to 25% of board thickness.

アレスト性が向上するメカニズムは、主き裂から分岐する分岐き裂が所定以上の長さで発生することで主き裂先端の応力緩和や閉口応力の発生により、破面形成エネルギーを消費や脆性破壊に対する駆動力の低下に起因して、主き裂の伝播を抑制するものと推定される。 The mechanism that improves the arrestability is that the fracture crack that branches off from the main crack occurs at a length longer than the specified length, thereby reducing stress at the tip of the main crack and generating a closing stress, thereby consuming fracture surface formation energy and brittleness. It is presumed that the propagation of the main crack is suppressed due to a decrease in the driving force against the fracture.

前記の様に、本発明が主き裂の傾斜角度θを１５°以上としたのは、これ未満であると前記分岐き裂の発生が困難になるためであって、これは破面形成エネルギーの消費が少なくなり主き裂の伝播阻止が困難になってアレスト性が低下するためである。
また、４５°以下としたのは、異方性が大きくなることを抑制するためである。 As described above, the reason why the inclination angle θ of the main crack is set to 15 ° or more in the present invention is that if it is less than this, the generation of the branched crack becomes difficult, and this is the fracture surface formation energy. This is because the consumption of slag decreases and it becomes difficult to prevent the propagation of the main crack, and the arrestability decreases.
The reason why the angle is 45 ° or less is to prevent anisotropy from increasing.

また、分岐き裂の発生長さを、板厚の５〜２５％としたのは、５％未満であると分岐き裂が形成されても、その分岐き裂が直ちに停止ことから主き裂の伝播を阻止する有効な抵抗になり難く、また板厚の２５％超では逆に主き裂と合体して、主き裂は抵抗をほとんど受けずに真直ぐ伝播してしまうことから、前記同様に極端にアレスト性が低下するためである。 In addition, the length of occurrence of the branch crack is 5 to 25% of the plate thickness. If it is less than 5%, even if the branch crack is formed, the branch crack immediately stops, so the main crack is stopped. Since it is difficult to become an effective resistance to prevent the propagation of cracks, and when it exceeds 25% of the plate thickness, it merges with the main crack, and the main crack propagates straight with almost no resistance. This is because the arrestability is extremely reduced.

図３は、アレスト靭性値Ｋｃａが―１０℃で５０００Ｎ／ｍｍ^1.5以上を示すアレスト性の優れた鋼板と、アレスト靭性値Ｋｃａが―１０℃で５０００Ｎ／ｍｍ^1.5未満のアレスト性の劣る鋼板（各鋼板の板厚：４０〜１００ｍｍ）を用いて前記同様の温度勾配型の標準ＥＳＳＯ試験を行い、その結果から、板厚表層部の（１００）Ｘ線面強度比と板厚中心部の（１１１）又は／及び（２１１）Ｘ線面強度比との関係を表したものである。なお、Ｘ線面強度比は、表層部に分岐き裂が発生した場合は分岐き裂発生領域の中心部と板厚中心部を測定し、分岐き裂が発生しなかった場合は表層から板厚の１５％の位置と板厚中心部を測定した。そして、このＸ線面強度比とは、Ｘ線回折法により求めた（１００）（１１１）（２１１）面回折強度のランダム方位試料の回折強度に対する相対比のことである。 Figure 3 is a steel sheet excellent in arrestability to arrest toughness Kca indicates 5000N / mm ^1.5 or more at -10 ° C., arrestability toughness Kca is poor arrestability of less than 5000N / mm ^1.5 at -10 ° C. steel (each The same temperature gradient type standard ESSO test was performed using a steel plate thickness of 40 to 100 mm. From the results, the (100) X-ray surface strength ratio of the plate thickness surface layer portion and the (111) of the plate thickness center portion were determined. ) Or / and (211) the relationship with the X-ray plane intensity ratio. Note that the X-ray surface strength ratio is determined by measuring the center of the branch crack generation region and the center of the plate thickness when a branch crack occurs in the surface layer, and from the surface layer when no branch crack occurs. The position of 15% of the thickness and the center of the plate thickness were measured. The X-ray plane intensity ratio is the relative ratio of the (100) (111) (211) plane diffraction intensity obtained by the X-ray diffraction method to the diffraction intensity of the random orientation sample.

この図３から、鋼板表裏面に、板厚の５〜２５％の長さの分岐き裂が発生するアレスト性の優れた鋼板は、板厚表層部の（１００）面強度比が１．５以上２．０未満、かつ、板厚中心部の（１１１）又は／及び（２１１）面強度比が２．０以上、５．０未満の範囲に分布することが判る。一方、鋼板表裏面に、前記分岐き裂が発生しないアレスト性の劣る鋼板は、上記以外の範囲に分布することが判る。 From FIG. 3, a steel plate with excellent arrestability in which a branch crack having a length of 5 to 25% of the plate thickness occurs on the front and back surfaces of the steel plate has a (100) plane strength ratio of the plate thickness surface layer portion of 1.5. It can be seen that the distribution is in the range of 2.0 or more and less than 5.0 with the (111) or / and (211) plane intensity ratio at the center of the plate thickness being less than 2.0. On the other hand, it can be seen that the inferior arrestability steel sheet in which the branch crack does not occur on the front and back surfaces of the steel sheet is distributed in a range other than the above.

つまり、板厚表層部の（１００）面強度比が１．５未満では分岐き裂がほとんど形成されないため有効な脆性破壊の抵抗になり難く、また２．０以上でも表裏層部に分岐き裂は発生せず逆に主き裂となり、脆性き裂は抵抗をほとんど受けずに真直ぐ伝播してしまうことから、表裏層部の（１００）面強度比を１．５以上２．０未満の範囲とした。さらに板厚中心部の（１１１）又は／及び（２１１）面強度比が２．０未満では主き裂は外部応力に垂直な面に１５°未満で傾斜して伝播してしまい、表裏層部での分岐き裂の発生が困難となり、表層部が有効な伝播抵抗とはならないことから、主き裂の傾斜角度を１５°以上にし、表層部に分岐き裂を発生させることによってアレスト性を向上させるために２．０以上とした。また、板厚中心部の当該集合組織強度を大きくし、傾斜角度が大きくなるほどアレスト性は向上するが、異方性が大きくなることを抑制するために、５．０未満にする。 That is, when the (100) plane strength ratio of the plate thickness surface layer portion is less than 1.5, almost no branch cracks are formed, so that it is difficult to provide effective resistance to brittle fracture. In contrast, the main crack does not occur, and the brittle crack propagates straight without receiving any resistance, so the (100) plane strength ratio of the front and back layer portions is in the range of 1.5 or more and less than 2.0. It was. Further, when the (111) and / or (211) plane strength ratio at the center of the plate thickness is less than 2.0, the main crack propagates at an angle of less than 15 ° to the plane perpendicular to the external stress, and the front and back layer portions Since it is difficult to generate a branch crack at the surface and the surface layer portion does not have an effective propagation resistance, the inclination angle of the main crack is set to 15 ° or more, and the arrest property is improved by generating a branch crack in the surface layer portion. In order to improve, it was set to 2.0 or more. In addition, the arrestability is improved as the texture strength of the central portion of the plate thickness is increased and the inclination angle is increased. However, in order to suppress anisotropy from increasing, it is set to less than 5.0.

この様に、板厚中心部における（１１１）又は／及び（２１１）面強度比が２．０以上と大きい領域では、へき開面が外部応力と垂直な面に対し傾斜しており、かつ集合組織強度が表裏層部のそれよりも大きいことから、主き裂８が発生し、これが外部応力と直角方向に対し傾斜して伝播する。このとき（100）面強度比が1．５より大きい表裏層部では、へき開面が外部応力と直角方向であり真っ直ぐ伝播し易いことから、主き裂８より分岐き裂５が発生する。しかし、分岐き裂５は、主き裂８に引っ張られ成長することはできず、さらに主き裂８が伝播することにより数ｍｍ以上の長さの分岐き裂が次々と形成されていくことになる。 In this way, in the region where the (111) or / and (211) plane strength ratio is as large as 2.0 or more at the center of the plate thickness, the cleavage plane is inclined with respect to the plane perpendicular to the external stress, and the texture Since the strength is greater than that of the front and back layer portions, a main crack 8 is generated, which propagates with an inclination to the direction perpendicular to the external stress. At this time, in the front and back layer portions where the (100) plane strength ratio is greater than 1.5, the cleaved surface is in a direction perpendicular to the external stress and easily propagates straight, so that the branch crack 5 is generated from the main crack 8. However, the branch crack 5 cannot be grown by being pulled by the main crack 8, and further, the branch crack having a length of several mm or more is formed one after another as the main crack 8 propagates. become.

上記のようなアレスト性向上効果は、降伏応力が３５５〜５５０ＭＰａである鋼板、及び板厚が４０〜１００ｍｍの鋼板において特に顕著となる。この理由は、降伏応力が３５５ＭＰａ未満又は５５０ＭＰａ超、板厚が４０ｍｍ未満又は１００ｍｍ超の領域では、本発明で規定しているような板厚方向に集合組織が異なるような分布を形成させ、脆性き裂の伝播挙動を制御させることが困難であるからである。 The effect of improving the arrestability as described above is particularly remarkable in a steel plate having a yield stress of 355 to 550 MPa and a steel plate having a thickness of 40 to 100 mm. The reason for this is that in the region where the yield stress is less than 355 MPa or more than 550 MPa and the plate thickness is less than 40 mm or more than 100 mm, a distribution in which the texture is different in the plate thickness direction as defined in the present invention is formed, and the brittleness This is because it is difficult to control the propagation behavior of the crack.

以下、各元素の量を限定した理由について説明する。
Ｃは厚手母材の強度を確保するために０．０４％以上必要であり、これが下限である。また、Ｃが０．１５％を超えると良好なＨＡＺ靭性を確保することが困難であることから、これが上限となるが、Ｃの上限は好ましくは０．１％である。 Hereinafter, the reason for limiting the amount of each element will be described.
C needs to be 0.04% or more in order to ensure the strength of the thick base material, and this is the lower limit. Further, if C exceeds 0.15%, it is difficult to ensure good HAZ toughness, so this is the upper limit, but the upper limit of C is preferably 0.1%.

Ｓｉは脱酸元素及び強化元素として有効であるため、０．１％以上必要であるが、０．５％を超えるとＨＡＺ靭性が著しく劣化するため、これが上限である。 Since Si is effective as a deoxidizing element and strengthening element, it needs to be 0.1% or more. However, if it exceeds 0.5%, the HAZ toughness deteriorates remarkably, so this is the upper limit.

Ｍｎは厚手母材の強度と靭性を経済的に確保するために０．５％以上必要である。ただし、２．５％を超えて添加すると、中心偏析が顕著となってこの部分の母材とＨＡＺ靭性が劣化するため、これが上限である。 Mn is required to be 0.5% or more in order to economically secure the strength and toughness of the thick base material. However, if added over 2.5%, the center segregation becomes prominent and this part of the base material and the HAZ toughness deteriorate, so this is the upper limit.

Ｐは不純物元素であり、靭性を安定的に確保するために０．０２％以下に低減する必要がある。 P is an impurity element and needs to be reduced to 0.02% or less in order to ensure toughness stably.

Ｓも不純物元素であり、Ｐと同様の理由で０．０１％以下に低減する必要がある。 S is also an impurity element and needs to be reduced to 0.01% or less for the same reason as P.

Ａｌは脱酸を担い、不純物元素であるＯを低減するために必要である。Ａｌ以外にもＳｉやＭｎも脱酸に寄与するが、たとえこれらの元素が添加される場合でも、０．００１％以上のＡｌがないと安定的にＯを抑えることは難しい。ただし、Ａｌが０．１％を超えると、アルミナ系の粗大酸化物やクラスターが生成し、母材とＨＡＺ靭性が劣化するため、これが上限である。 Al is necessary for deoxidizing and reducing O which is an impurity element. In addition to Al, Si and Mn also contribute to deoxidation, but even when these elements are added, it is difficult to stably suppress O without 0.001% or more of Al. However, if Al exceeds 0.1%, an alumina-based coarse oxide or cluster is generated, and the base material and the HAZ toughness deteriorate, so this is the upper limit.

ＴｉはＴｉＮを形成することによって、鋼片加熱時や溶接時にオーステナイト粒径が大きくなることを抑制でき、母材とＨＡＺ靭性を向上させる効果がある。この効果を得るためには０．００５％以上必要である。しかし、過剰なＴｉの添加は、ＴｉＣ形成によりＨＡＺ靭性が劣化するため、０．０２％を上限とする。 By forming TiN, Ti can suppress an increase in the austenite grain size at the time of steel slab heating or welding, and has the effect of improving the base material and the HAZ toughness. In order to obtain this effect, 0.005% or more is necessary. However, excessive Ti addition causes the HAZ toughness to deteriorate due to TiC formation, so 0.02% is made the upper limit.

Ｎは上記したようにＴｉＮ形成による母材とＨＡＺ靭性向上効果を得るために０．００１％以上必要である。しかし、過剰なＮの添加は鋳片割れや母材とＨＡＺ靭性の劣化を招くため、０．００８％を上限とする。 As described above, N is required to be 0.001% or more in order to obtain the base material and the HAZ toughness improvement effect by TiN formation. However, excessive addition of N causes cracks in the slab and deterioration of the base metal and HAZ toughness, so 0.008% is made the upper limit.

また、上記した添加元素の他に、質量％で、Ｃｕ：０．０５〜１％、Ｎｉ：０．０５〜２％、Ｃｒ：０．０５〜１％、Ｍｏ：０．０５〜０．５％、Ｎｂ：０．００３〜０．１％、Ｖ：０．００５〜０．２％、Ｂ：０．０００２〜０．００３％の少なくとも１種以上を化学成分として含有しても良い。これらを添加することにより、母材の強度と靭性が確保される。ただし、これらの元素が多すぎると母材とＨＡＺ靭性や溶接性が低下するため、それぞれの元素に上限を設ける必要がある。 In addition to the additive elements described above, in terms of mass%, Cu: 0.05 to 1%, Ni: 0.05 to 2%, Cr: 0.05 to 1%, Mo: 0.05 to 0.5 %, Nb: 0.003 to 0.1%, V: 0.005 to 0.2%, B: 0.0002 to 0.003% may be contained as a chemical component. By adding these, the strength and toughness of the base material are ensured. However, if there are too many of these elements, the base material and the HAZ toughness and weldability will deteriorate, so it is necessary to provide an upper limit for each element.

さらに、上記した添加元素の他に、質量％で、Ｃａ：０．０００３〜０．００５％、Ｍｇ：０．０００３〜０．００５％、ＲＥＭ：０．０００３〜０．００５％の少なくとも１種以上を化学成分として含有しても良い。これらを添加することにより、酸化物、硫化物がピン止め粒子となりオーステナイト粒の成長を抑制する、またはフェライト変態核となりＨＡＺ組織を微細化することによりＨＡＺ靭性が向上する。ただし、これらの元素が多すぎると粗大な介在物、クラスターが形成され、ＨＡＺ靭性が劣化するため、それぞれの元素に上限を設ける必要がある。 Further, in addition to the additive elements described above, at least one of Ca: 0.0003 to 0.005%, Mg: 0.0003 to 0.005%, and REM: 0.0003 to 0.005% in mass%. You may contain the above as a chemical component. By adding these, oxides and sulfides become pinned particles to suppress the growth of austenite grains, or ferrite transformation nuclei to refine the HAZ structure by improving the HAZ toughness. However, if there are too many of these elements, coarse inclusions and clusters are formed and the HAZ toughness deteriorates, so it is necessary to provide an upper limit for each element.

以下、本発明の製造方法を限定した理由について説明する。 Hereinafter, the reason why the production method of the present invention is limited will be described.

まず、上記した適切な化学成分組成に調整した溶鋼を、転炉等の通常公知の溶製方法で溶製し、連続鋳造等の通常公知の鋳造方法で鋼素材とする。 First, the molten steel adjusted to the appropriate chemical composition described above is melted by a generally known melting method such as a converter, and is made into a steel material by a generally known casting method such as continuous casting.

次に、鋼素材を９５０℃〜１２５０℃の温度に加熱し、オーステナイト単相化する。これは９５０℃未満ではオーステナイト単相化が不十分であり、１２５０℃超では加熱γ粒径が極端に粗大化して圧延後に微細な組織を得ることが困難となり靭性が低下するからである。この加熱した鋼素材は、オーステナイトの細粒化を目的に９００℃以上での再結晶圧延を行っても良いが、圧延なしのままでも構わない。 Next, the steel material is heated to a temperature of 950 ° C. to 1250 ° C. to make an austenite single phase. This is because if the temperature is lower than 950 ° C., the austenite single phase is insufficient, and if it exceeds 1250 ° C., the heated γ grain size becomes extremely coarse and it becomes difficult to obtain a fine structure after rolling, and the toughness decreases. This heated steel material may be recrystallized and rolled at 900 ° C. or higher for the purpose of refining austenite, but it may be left unrolled.

引き続き行う圧延の過程が本発明の最も重要な部分である。すなわち、表面温度を６５０℃以上８５０℃以下、板厚中心温度を８５０℃超１０５０℃以下で累積圧下率３０％以上の圧延を行うことが必要である。 The subsequent rolling process is the most important part of the present invention. That is, it is necessary to perform rolling with a surface temperature of 650 ° C. or more and 850 ° C. or less, a sheet thickness center temperature of more than 850 ° C. and 1050 ° C. or less, and a cumulative reduction ratio of 30% or more.

まず、表面温度範囲を６５０〜８５０℃とした理由は、分岐き裂が形成される表裏層部において、圧延面と平行な（１００）集合組織を発達させるためには未再結晶域オーステナイト、あるいはオーステナイト／フェライト二相域で仕上げ圧延する必要があるからである。６５０℃未満では（１００）面強度比が２．０以上となり、８５０℃超では逆に（１００）面強度比が１．５未満となるために分岐き裂が形成されずアレスト性向上が図れないのでこの範囲とした。母材強度、靭性、生産性とのバランスを考慮すると、７００〜８００℃とするのが望ましい。 First, the reason why the surface temperature range is set to 650 to 850 ° C. is that, in order to develop a (100) texture parallel to the rolling surface in the front and back layer portions where branch cracks are formed, This is because it is necessary to finish-roll in the austenite / ferrite two-phase region. If it is less than 650 ° C., the (100) plane strength ratio is 2.0 or more, and if it exceeds 850 ° C., on the contrary, the (100) plane strength ratio is less than 1.5. Because there is no, this range. Considering the balance between the base material strength, toughness, and productivity, the temperature is preferably 700 to 800 ° C.

次に、板厚中心部計算温度を８５０℃超１０５０℃以下とした理由は、分岐き裂が形成される表裏層部以外の領域において、圧延面と平行な（１１１）又は／及び（２１１）集合組織を発達させるためにはオーステナイト単相域、少なくとも再結晶オーステナイト域で圧延する必要があるからである。８５０℃以下では、（１１１）又は／及び（２１１）面強度比が２．０未満となり、主き裂を傾斜させることができないので、これを下限とする。また１０５０℃超では（１１１）又は／及び（２１１）面強度比が２．０以上を確保できるものの組織粗大化により母材靭性が劣化してしまうので、これを上限とする。 Next, the reason why the calculation temperature at the center of the plate thickness is more than 850 ° C. and 1050 ° C. or less is that (111) or / and (211) parallel to the rolling surface in the region other than the front and back layer portions where the branch crack is formed. This is because rolling in the austenite single phase region, at least the recrystallized austenite region, is necessary to develop the texture. At 850 ° C. or lower, the (111) or / and (211) plane strength ratio is less than 2.0, and the main crack cannot be inclined, so this is the lower limit. Further, when the temperature exceeds 1050 ° C., the (111) or / and (211) plane strength ratio can be 2.0 or more, but the base material toughness deteriorates due to the coarsening of the structure, so this is the upper limit.

上記のように板厚方向に集合組織分布を形成させるためには、板厚方向の温度偏差を大きくすることが重要である。このような温度偏差を容易に実現させるためには、板面の冷却が不均一になりにくい1〜１０℃／ｓ程度の冷却速度で圧延前に水冷を行うと良い。 In order to form a texture distribution in the thickness direction as described above, it is important to increase the temperature deviation in the thickness direction. In order to easily realize such a temperature deviation, it is preferable to perform water cooling before rolling at a cooling rate of about 1 to 10 ° C./s in which the cooling of the plate surface is difficult to be uneven.

次に、圧延の累積圧下率を３０％以上とする理由は、３０％未満では温度条件を満足しても集合組織の発達や靭性の確保が困難であるので、これを下限とする。集合組織の発達や強度、靭性、生産性バランスの観点からは、４５〜７０％とするのが望ましい。 Next, the reason why the cumulative rolling reduction ratio is set to 30% or more is that if it is less than 30%, it is difficult to secure the development of texture and toughness even if the temperature condition is satisfied. From the viewpoint of texture development, strength, toughness, and productivity balance, it is desirable to set it to 45 to 70%.

上記の圧延後、６００℃以上から、板厚平均で８℃/s以上の冷却速度で５００℃以下の温度まで加速冷却を行う必要がある。 After the rolling described above, it is necessary to perform accelerated cooling from 600 ° C. or higher to a temperature of 500 ° C. or lower at a cooling rate of 8 ° C./s or higher in average thickness.

冷却開始温度を６００℃以上とした理由は、６００℃未満では冷却前にフェライト変態が進行し、フェライト主体の組織となってしまい、強度の確保が困難になるからである。 The reason for setting the cooling start temperature to 600 ° C. or higher is that if it is less than 600 ° C., the ferrite transformation proceeds before cooling, resulting in a structure mainly composed of ferrite, making it difficult to ensure the strength.

加速冷却時の冷却速度を板厚平均で８℃／ｓ以上とした理由は、上記と同様の理由で８℃／ｓ未満ではベイナイト主体組織が得られず、強度の確保が困難になるからである。また、分岐き裂が形成される表裏層部以外の領域では、ベイナイト主体組織でなければ（１１１）又は／及び（２１１）集合組織を発達させることができないので、これを下限とする。 The reason why the cooling rate at the time of accelerated cooling is 8 ° C./s or more on average in the plate thickness is that, for the same reason as described above, if it is less than 8 ° C./s, a bainite main structure cannot be obtained and it is difficult to ensure strength. is there. Further, in a region other than the front and back layer portions where the branch crack is formed, the (111) or / and (211) texture cannot be developed unless it is a bainite-based structure, and this is set as the lower limit.

５００℃以下の温度まで加速冷却する理由は、上記と同様の理由で５００℃超ではベイナイト主体組織が得られず強度確保が困難である上に、アレスト性向上に必要な集合組織要件を満足できないからである。 The reason for accelerated cooling to a temperature of 500 ° C. or lower is that, for the same reason as above, if the temperature exceeds 500 ° C., a bainite-based structure cannot be obtained and it is difficult to ensure strength, and the texture requirement necessary for improving arrestability cannot be satisfied. Because.

加速冷却後、強度と靭性を調整する目的で必要に応じ３００〜６５０℃の温度で焼き戻しすることが可能である。その効果を得るためには３００℃以上にする必要があり、６５０℃超では極端に軟化し強度の確保が困難となるので、６５０℃を上限とする。 After accelerated cooling, it can be tempered at a temperature of 300 to 650 ° C. as necessary for the purpose of adjusting strength and toughness. In order to obtain the effect, it is necessary to set the temperature to 300 ° C. or higher. If it exceeds 650 ° C., it becomes extremely soft and it is difficult to ensure the strength.

このように、極低温圧延、及び複雑な熱処理工程を必要としないため、本実施形態に係る高強度厚鋼板は生産性が高く、かつ低コストになる。また、残留応力も抑制されるため、形状矯正に起因したコスト増加を抑制できる。 Thus, since cryogenic rolling and a complicated heat treatment process are not required, the high-strength thick steel plate according to this embodiment has high productivity and low cost. Moreover, since residual stress is also suppressed, the cost increase resulting from shape correction can be suppressed.

以上のように本実施形態によれば、化学成分、製造条件を適切な値に制御し、かつ鋼板の板厚方向の集合組織強度と分布を制御することによって、脆性き裂の伝播方向や破面形態を制御することができる。これにより、高強度厚鋼板において、アレスト性を向上させることができる。そして、降伏応力が３５５〜５５０ＭＰａ、かつ板厚が４０〜１００ｍｍの鋼板において、−１０℃のアレスト靭性値であるＫｃａを５０００Ｎ／ｍｍ^1.5以上にすることができる。また、極端な低温圧延、及び複雑な熱処理工程を必要としないため生産性が高く、低コストにすることができる。 As described above, according to the present embodiment, by controlling the chemical composition and manufacturing conditions to appropriate values and controlling the texture strength and distribution in the thickness direction of the steel sheet, the propagation direction and fracture of the brittle crack are controlled. The surface form can be controlled. Thereby, arrestability can be improved in a high-strength thick steel plate. The yield stress 355～550MPa, and plate thickness in the steel sheet of 40 to 100 mm, can be a Kca a arrestability toughness of -10 ° C. to 5000N / mm ^1.5 or more. In addition, since extremely low temperature rolling and complicated heat treatment steps are not required, productivity is high and costs can be reduced.

製鋼工程において溶鋼の化学成分調整を行った後、連続鋳造によって鋳片を製造した。表１に化学成分を示す。この鋳片を用いて板厚４０〜１００ｍｍの厚鋼板を製造した。表２に各厚鋼板の製造方法示す。なお、板厚中心温度は、通常公知の差分法による熱伝導解析により求めた。表２には板厚中心部のミクロ組織構成も合わせて示す。 After adjusting the chemical composition of the molten steel in the steel making process, a slab was produced by continuous casting. Table 1 shows chemical components. A thick steel plate having a thickness of 40 to 100 mm was manufactured using this slab. Table 2 shows a method for manufacturing each thick steel plate. The plate thickness center temperature was obtained by heat conduction analysis by a generally known differential method. Table 2 also shows the microstructure of the center of the plate thickness.

各厚鋼板の集合組織強度、機械的特性（引張特性、衝撃特性）、アレスト性、及び脆性破面形態（マクロなき裂、すなわち主き裂の外部応力と垂直な面に対する傾斜角度、及び表層部の分岐き裂領域の板厚方向長さ）を測定した。集合組織強度はＸ線回折法により、（１００）（１１１）（２１１）面回折強度のランダム方位試料の回折強度に対する相対比を求めた。表層部は板厚の５、１５、２５％の位置の（１００）面強度比、そして板厚中心部の（１１１）又は（２１１）面強度比をそれぞれ記載した。引張特性は、ＪＩＳＺ２２４１に準拠し、ＪＩＳＺ２２０１の丸棒引張試験片を板厚中心部から圧延方向と直角方向に２本採取し引張試験に供し、降伏応力（ＹＰ）及び引張強さ（ＴＳ）のそれぞれの平均値を記載した。また衝撃特性は、ＪＩＳＺ２２４２に準拠し、ＪＩＳＺ２２０２のＶノッチシャルピー衝撃試験片を板厚中心部から圧延方向と平行に３本採取しシャルピー衝撃試験に供し、−４０℃でのシャルピー吸収エネルギー（ｖＥ−４０）の平均値を記載した。アレスト性は、板幅５００ｍｍの温度勾配型ＥＳＳＯ試験により、−１０℃でのアレスト靭性値Ｋｃａを求めた。脆性破面形態は、ＥＳＳＯ試験後の破面を用いて、マクロなき裂、すなわち主き裂の外部応力と垂直な面に対する傾斜角度、及び表層部の分岐き裂領域の板厚方向長さを測定した。表３に測定結果を示す。 Texture strength, mechanical properties (tensile properties, impact properties), arrestability, and brittle fracture surface morphology of each steel plate (macro crack, that is, the inclination angle of the main crack to the plane perpendicular to the external stress, and surface layer part The thickness in the thickness direction of the branched crack region was measured. For the texture intensity, the relative ratio of the (100) (111) (211) plane diffraction intensity to the diffraction intensity of the randomly oriented sample was determined by X-ray diffraction. For the surface layer portion, the (100) plane strength ratio at the positions of 5, 15 and 25% of the plate thickness and the (111) or (211) plane strength ratio at the plate thickness center portion are described. Tensile properties are based on JIS Z 2241. Two round bars of JIS Z 2201 are taken from the center of the plate thickness in the direction perpendicular to the rolling direction and subjected to a tensile test, yield stress (YP) and tensile strength. Each average value of (TS) was described. In addition, the impact characteristics are based on JIS Z 2242. Three V-notch Charpy impact test pieces of JIS Z 2202 are sampled from the center of the plate thickness in parallel with the rolling direction and subjected to Charpy impact test. Charpy absorption at -40 ° C The average value of energy (vE-40) was described. For the arrestability, an arrest toughness value Kca at −10 ° C. was obtained by a temperature gradient type ESSO test with a plate width of 500 mm. The brittle fracture surface morphology uses the fracture surface after the ESSO test to determine the macro crack, that is, the inclination angle of the main crack to the plane perpendicular to the external stress and the length in the plate thickness direction of the branch crack area in the surface layer. It was measured. Table 3 shows the measurement results.

鋼番１〜６は本発明の厚鋼板である。化学成分、製造方法ともに本発明範囲内であるため、集合組織強度、分布、脆性破面形態も本発明要件を満足している。この結果、−１０℃でのアレスト靭性値Ｋｃａが５０００Ｎ／ｍｍ^1.5以上の優れた値を示していた。また、機械的性質も、降伏強度（ＹＰ）が３６０〜５４０ＭＰａ、引張強さ（ＴＳ）が５００〜６８０ＭＰａの高強度、−４０℃シャルピー吸収エネルギー（ｖＥ−４０）が１５０〜３００Ｊの高靭性を示していた。 Steel numbers 1 to 6 are the thick steel plates of the present invention. Since both the chemical composition and the production method are within the scope of the present invention, the texture strength, distribution, and brittle fracture surface form also satisfy the requirements of the present invention. As a result, the arrest toughness value Kca at −10 ° C. was an excellent value of 5000 N / mm ^1.5 or more. In addition, the mechanical properties include high strength with yield strength (YP) of 360-540 MPa, tensile strength (TS) of 500-680 MPa, and high toughness of −40 ° C. Charpy absorbed energy (vE-40) of 150-300 J. Was showing.

これに対し、鋼番７〜１３は比較例となる厚鋼板である。鋼番７、８、９、１０は製造方法要件のうち、仕上圧延開始時の表面温度と板厚中心温度のいずれか、または両方が満足していない。そのうち、鋼番７は、仕上げ圧延前の表面温度が本発明で規定した下限を下回る６００℃と低いため、ミクロ組織がフェライト主体となり母材強度や靭性が低い値を示し、かつ、表層部の集合組織強度が本発明の上限を上回り、中心部の集合組織強度が本発明の下限を下回り、脆性き裂は真直ぐ伝播し、分岐き裂も生じなかったため、Ｋｃａが３０００Ｎ／ｍｍ^1.5と極めて低い値を示した。また、鋼番８は、仕上圧延開始時の板厚中心（ｔ／２開始）温度が本発明で規定した下限を下回る７８０℃であり、そのため板厚中心部の集合組織強度が本発明の下限を下回り、脆性き裂は真直ぐ伝播し、分岐き裂も生じなかったことから、アレスト性の大幅な改善は認められなかった。また、鋼番９は、仕上圧延開始時の板厚中心温度が本発明の上限を上回る１１００℃であり、そのため、母材の靭性が極端に低く溶接構造用鋼としての使用性能を満足していない。表層部の集合組織要件も満足しておらず、主き裂は傾斜して伝播するものの、表層部の分岐き裂領域の板厚方向長さが２ｍｍであり、板厚の３％であったため、本発明の下限を外れ脆性破壊の抵抗とはなり難く、Ｋｃａは３８００Ｎ／ｍｍ^1.5と低い。また、鋼番１０は、仕上圧延開始時の表面温度と板厚中心温度ともに本発明の上限を外れている。そのため母材靭性の著しい劣化があり、集合組織要件も満足しておらず、マクロなき裂は傾斜するものの分岐き裂が形成されないことから、Ｋｃａは３０００Ｎ／ｍｍ^1.5と極めて低い値を示した。 On the other hand, steel numbers 7 to 13 are thick steel plates as comparative examples. Steel Nos. 7, 8, 9, and 10 do not satisfy either or both of the surface temperature and the plate thickness center temperature at the start of finish rolling, among the manufacturing method requirements. Among them, steel No. 7 has a surface temperature before finish rolling as low as 600 ° C. which is lower than the lower limit defined in the present invention. Therefore, the microstructure is mainly composed of ferrite and exhibits a low base metal strength and toughness, and the surface layer portion The texture strength exceeds the upper limit of the present invention, the texture strength of the central portion is lower than the lower limit of the present invention, and the brittle crack propagates straight and no branch crack occurs. Therefore, Kca is as low as 3000 N / mm ^1.5. The value is shown. Steel No. 8 has a sheet thickness center (t / 2 start) temperature at the start of finish rolling of 780 ° C. which is lower than the lower limit defined in the present invention. Therefore, the texture strength at the center of the sheet thickness is the lower limit of the present invention. The brittle crack propagated straight and the branch crack did not occur, so there was no significant improvement in arrestability. Steel No. 9 has a plate thickness center temperature at the start of finish rolling of 1100 ° C., which exceeds the upper limit of the present invention, and therefore the toughness of the base metal is extremely low and satisfies the use performance as a welded structural steel. Absent. Although the texture requirement of the surface layer portion is not satisfied and the main crack propagates in an inclined manner, the length in the plate thickness direction of the branch crack region of the surface layer portion is 2 mm, which is 3% of the plate thickness The lower limit of the present invention is not exceeded, and resistance to brittle fracture is difficult, and Kca is as low as 3800 N / mm ^1.5 . Steel No. 10 is outside the upper limit of the present invention in both the surface temperature and the plate thickness center temperature at the start of finish rolling. For this reason, the toughness of the base metal is remarkably deteriorated, the texture requirements are not satisfied, and the macro crack is inclined but the branch crack is not formed. Therefore, Kca showed an extremely low value of 3000 N / mm ^1.5 .

また、鋼番１１は仕上圧延での累積圧下率（ＣＲ率）が本発明の下限を下回り、そのため表層部の集合組織強度が本発明の下限を下回り、マクロなき裂は若干傾斜するものの脆性破壊の抵抗となる分岐き裂が形成されず、アレスト性は低下した。 Steel No. 11 has a cumulative rolling reduction (CR ratio) in finish rolling that is lower than the lower limit of the present invention. Therefore, the texture strength of the surface layer is lower than the lower limit of the present invention, and the macro crack is slightly inclined but brittle fracture occurs. As a result, no branch crack was formed, and the arrestability decreased.

また、鋼番１２、１３は仕上圧延後の冷却条件を満足していない。そのうち、鋼番１２は、冷却終了温度が本発明の下限を上回る５５０℃であり、鋼番１３は空冷している。そのため、両鋼ともにミクロ組織はフェライト主体となり母材の強度が低い。さらに表層部の集合組織要件は満足しているものの、板厚中心部の（１１１）又は／及び（２１１）面強度比が低く本発明の下限を下回るため、脆性き裂は真直ぐ伝播し、分岐き裂も生じず、Ｋｃａは低い値を示した。 Steel numbers 12 and 13 do not satisfy the cooling conditions after finish rolling. Among them, Steel No. 12 has a cooling end temperature of 550 ° C. exceeding the lower limit of the present invention, and Steel No. 13 is air-cooled. Therefore, the microstructure of both steels is mainly ferrite, and the strength of the base material is low. Furthermore, although the texture requirement of the surface layer is satisfied, the (111) or / and (211) plane strength ratio at the center of the plate thickness is low and below the lower limit of the present invention, so that the brittle crack propagates straight and branches. No cracks occurred and Kca showed a low value.

以上の実施例から、本発明を適用することにより、降伏応力が３５５〜５５０ＭＰａ、板厚が４０〜１００ｍｍ、かつ−１０℃のアレスト靭性値Ｋｃａが５０００Ｎ／ｍｍ^1.5以上である脆性き裂伝播停止性能に優れた高強度厚鋼板及びその製造方法を提供することが確認された。 From the above examples, by applying the present invention, a brittle crack propagation stop with a yield stress of 355 to 550 MPa, a plate thickness of 40 to 100 mm, and an arrest toughness value Kca of −10 ° C. of 5000 N / mm ^1.5 or more. It has been confirmed that a high-strength thick steel plate excellent in performance and a method for producing the same are provided.

なお、本発明は上述した実施形態に限定されるものではなく、本発明の主旨を逸脱しない範囲内で種々変更して実施することが可能である。 Note that the present invention is not limited to the above-described embodiment, and various modifications can be made without departing from the spirit of the present invention.

温度勾配型標準ＥＳＳＯ試験を行った後の脆性き裂の伝播挙動及び脆性破面を示す簡略図である。It is a simplified diagram showing a brittle crack propagation behavior and a brittle fracture surface after performing a temperature gradient type standard ESSO test. 種々の鋼板に発生する主き裂の傾斜角度と分岐き裂寸法の関係を示す図である。It is a figure which shows the relationship between the inclination angle of the main crack which generate | occur | produces in various steel plates, and a branch crack size. 種々の鋼板の板厚表層部の（１００）Ｘ線面強度比と板厚中心部の（１１１）又は／及び（２１１）Ｘ線面強度比との関係を示す図である。It is a figure which shows the relationship between the (100) X-ray surface intensity ratio of the plate | board thickness surface layer part of various steel plates, and the (111) or / and (211) X-ray plane intensity ratio of a plate | board thickness center part.

符号の説明Explanation of symbols

１外部応力
２予め形成した切欠き
３マクロな脆性き裂
４マクロき裂の伝播方向
５分岐き裂
６外部応力と直角方向
７鋼板の表裏層部
８主き裂
９板厚方向中心部 DESCRIPTION OF SYMBOLS 1 External stress 2 Preformed notch 3 Macro brittle crack 4 Macro crack propagation direction 5 Branch crack 6 Direction perpendicular to external stress 7 Front and back layer part of steel sheet 8 Main crack 9 Thickness direction center part

Claims

質量％で、
Ｃ：０．０４〜０．１５％、
Ｓｉ：０．１〜０．５％、
Ｍｎ：０．５〜２．５％、
Ｐ：≦０．０２％、
Ｓ：≦０．０１％、
Ａｌ：０．００１〜０．１％、
Ｔｉ：０．００５〜０．０２％、
Ｎ：０．００１〜０．００８％
を含有し、残部が鉄及び不可避不純物によって化学成分が構成された鋼板で、ミクロ組織がベイナイトを母相としたフェライト又は／及びパーライト組織であり、鋼板の表面及び裏面から板厚の２５％までの表裏層部とそれ以外の板厚中心部の三層に分けたとき、表裏層部で、板厚の５％以上２５％以下の厚さ領域に圧延面と平行な（１００）Ｘ線面強度比が１．５以上２．０未満の集合組織を有しており、それ以外の板厚中心部を含む領域において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比が２．０以上の集合組織を有していることを特徴とする脆性き裂伝播停止性能に優れた高強度厚鋼板。 % By mass
C: 0.04 to 0.15%,
Si: 0.1 to 0.5%,
Mn: 0.5 to 2.5%
P: ≦ 0.02%,
S: ≦ 0.01%
Al: 0.001 to 0.1%,
Ti: 0.005 to 0.02%,
N: 0.001 to 0.008%
The balance is a steel plate in which the chemical composition is composed of iron and inevitable impurities, and the microstructure is a ferrite or / and pearlite structure with bainite as a parent phase, from the front and back surfaces of the steel plate to 25% of the plate thickness (100) X-ray plane parallel to the rolling surface in the thickness region of 5% or more and 25% or less of the plate thickness at the front and back layer portions when divided into three layers of the front and back layer portions and the other thickness center portion (111) or / and (211) X-ray plane intensity ratio parallel to the rolling surface in the region including the central portion of the plate thickness that has a texture with an intensity ratio of 1.5 or more and less than 2.0 Is a high-strength thick steel plate excellent in brittle crack propagation stopping performance, characterized by having a texture of 2.0 or more.

質量％で、
Ｃｕ：０．０５〜１％、
Ｎｉ：０．０５〜２％、
Ｃｒ：０．０５〜１％、
Ｍｏ：０．０５〜０．５％、
Ｎｂ：０．００３〜０．１％、
Ｖ：０．００５〜０．２％、
Ｂ：０．０００２〜０．００３％
の少なくとも１種以上を化学成分として含有することを特徴とする請求項１に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 % By mass
Cu: 0.05 to 1%,
Ni: 0.05-2%,
Cr: 0.05 to 1%,
Mo: 0.05-0.5%
Nb: 0.003 to 0.1%,
V: 0.005 to 0.2%,
B: 0.0002 to 0.003%
The high-strength thick steel plate having excellent brittle crack propagation stopping performance according to claim 1, comprising at least one of the above as chemical components.

質量％で、
Ｃａ：０．０００３〜０．００５％、
Ｍｇ：０．０００３〜０．００５％、
ＲＥＭ：０．０００３〜０．００５％
の少なくとも１種以上を化学成分として含有することを特徴とする請求項１または２に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 % By mass
Ca: 0.0003 to 0.005%,
Mg: 0.0003 to 0.005%,
REM: 0.0003 to 0.005%
The high-strength thick steel plate excellent in brittle crack propagation stopping performance according to claim 1 or 2, wherein at least one of the above is contained as a chemical component.

板厚が４０ｍｍ以上であることを特徴とする請求項１〜３のいずれか１項に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 The high-strength thick steel plate excellent in brittle crack propagation stopping performance according to any one of claims 1 to 3, wherein the plate thickness is 40 mm or more.

前記表裏層部が分岐き裂となりに、そのき裂が板厚の５％以上２５％以下の長さで外部応力と垂直方向に伝播し、それ以外の板厚中心部の領域が主き裂となり外部応力と垂直な面に対し１５°以上４５°以下で伝播することを特徴とする請求項１〜４のいずれかに記載の脆性き裂伝播停止性能に優れた高強度厚鋼板。 The front and back layer portions become branched cracks, and the crack propagates in a direction perpendicular to the external stress with a length of 5% or more and 25% or less of the plate thickness, and the other region in the center portion of the plate thickness is the main crack. The high-strength thick steel plate excellent in brittle crack propagation stopping performance according to any one of claims 1 to 4, wherein the high-strength steel plate is propagated at 15 ° or more and 45 ° or less with respect to a plane perpendicular to external stress.

請求項１〜３のいずれかに記載の化学成分を有する鋼片を、９５０〜１２５０℃に加熱し、表面温度を６５０℃以上８５０℃以下、板厚中心温度を８５０℃超１０５０℃以下で累積圧下率３０％以上の圧延を行った後、表面温度６００℃以上から、板厚平均で８℃/s以上の冷却速度で５００℃以下の温度まで加速冷却を行うことを特徴とする脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 The steel slab having the chemical component according to any one of claims 1 to 3 is heated to 950 to 1250 ° C, the surface temperature is 650 ° C or higher and 850 ° C or lower, and the sheet thickness center temperature is accumulated above 850 ° C and below 1050 ° C. A brittle crack characterized by performing accelerated cooling from a surface temperature of 600 ° C. or more to a temperature of 500 ° C. or less at a cooling rate of 8 ° C./s or more on the average sheet thickness after rolling at a reduction rate of 30% or more. A method for producing high-strength thick steel plates with excellent propagation stop performance.

前記加速冷却を終了した後、３００℃以上６５０℃以下で焼戻しすることを特徴とする請求項６に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 The method for producing a high-strength thick steel plate excellent in brittle crack propagation stopping performance according to claim 6, wherein after accelerating cooling is finished, tempering is performed at 300 ° C. or more and 650 ° C. or less.

板厚が４０ｍｍ以上であることを特徴とする請求項６または７に記載の脆性き裂伝播停止性能に優れた高強度厚鋼板の製造方法。 The method for producing a high-strength thick steel plate excellent in brittle crack propagation stopping performance according to claim 6 or 7, wherein the plate thickness is 40 mm or more.