JP6274375B1 - High strength thick steel plate and manufacturing method thereof - Google Patents
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Abstract
大入熱溶接部の靭性および脆性亀裂伝播停止特性に優れた高強度厚鋼板およびその製造方法を提供する。所定の成分組成を有し、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比が１．２以上、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比が１．５以上かつ（２２２）面Ｘ線強度比が２．５以下であり、３００ｋＪ／ｃｍ超の入熱量で溶接した溶接継手ボンド部の−４０℃におけるシャルピー吸収エネルギー（ｖＥ−４０）が８０Ｊ以上であり、ＥＳＳＯ試験による−１０℃におけるＫｃａ値（Ｋｃａ（−１０℃））が６０００Ｎ／ｍｍ３／２以上であり、引張り強さが５８０ＭＰａ以上であり、板厚が５０ｍｍ超である、高強度厚鋼板。所定の成分組成を有する鋼素材を、１０００〜１２００℃の温度に加熱したのち、板厚中央が、オーステナイト再結晶温度域、次いで、オーステナイト未再結晶温度域で熱間圧延を行い、前記熱間圧延の途中で、鋼板の表裏面から加熱を行い、少なくとも前記オーステナイト未再結晶温度域での熱間圧延の一部を、鋼板表面の温度−板厚中央の温度≧−４０℃となる条件で行うように板厚方向の温度分布を制御し、熱間圧延終了時における鋼板表面と板厚中央の温度差を５℃以内とし、かつ、板厚中央の温度をＡｒ３℃〜（Ａｒ３+３０）℃とする、高強度厚鋼板の製造方法。Provided are a high-strength thick steel plate excellent in toughness and brittle crack propagation stopping characteristics of a high heat input weld and a method for producing the same. The (211) plane X-ray intensity ratio in the plane parallel to the plate surface of the steel sheet surface having a predetermined component composition is 1.2 or more in the plane parallel to the plate surface of the steel plate surface, Charpy absorbed energy (vE-40) at −40 ° C. of a welded joint portion having a ratio of 1.5 or more and a (222) plane X-ray intensity ratio of 2.5 or less and welded with a heat input of more than 300 kJ / cm Is 80 J or more, the Kca value (Kca (−10 ° C.)) at −10 ° C. by the ESSO test is 6000 N / mm 3/2 or more, the tensile strength is 580 MPa or more, and the plate thickness is more than 50 mm. High strength thick steel plate. After heating the steel material having a predetermined component composition to a temperature of 1000 to 1200 ° C., the center of the plate thickness is subjected to hot rolling in the austenite recrystallization temperature range, and then in the austenite non-recrystallization temperature range, During the rolling, heating is performed from the front and back surfaces of the steel sheet, and at least a part of the hot rolling in the austenite non-recrystallization temperature range is performed under the condition that the temperature of the steel sheet surface-the temperature at the center of the sheet thickness is -40C. The temperature distribution in the plate thickness direction is controlled so that the temperature difference between the steel plate surface and the plate thickness center at the end of hot rolling is within 5 ° C., and the plate thickness center temperature is Ar 3 ° C. to (Ar 3 + 30) A method for producing a high-strength thick steel plate at ℃.
Description
本発明は、高強度厚鋼板およびその製造方法に関する。本発明は、特に、船舶、海洋構造物、低温貯蔵タンク、建築・土木構造物等の大型構造物に使用する大入熱溶接部の靭性および脆性亀裂伝播停止特性に優れた高強度厚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. In particular, the present invention is a high-strength thick steel plate excellent in toughness and brittle crack propagation stopping properties of large heat input welds used for large structures such as ships, offshore structures, low-temperature storage tanks, and construction / civil engineering structures, and It relates to the manufacturing method.
船舶や、海洋構造物、低温貯蔵タンク、建築・土木構造物等の大型構造物においては、脆性破壊に伴う事故が起きると、社会経済や環境などに及ぼす影響が大きいため、安全性の向上が常に求められ、使用される鋼材に対しては、使用温度における靭性や、脆性亀裂伝播停止特性が高いレベルで要求されている。 For large structures such as ships, marine structures, low-temperature storage tanks, and construction / civil engineering structures, accidents associated with brittle fractures have a large impact on the socio-economic environment and the environment. Steel materials that are always required and used are required to have high levels of toughness at operating temperature and brittle crack propagation stopping characteristics.
たとえば、コンテナ船やバルクキャリアーなどの船舶は、その構造上、船体外板に高強度の厚肉材を使用するが、最近では、船体の大型化に伴って一層の高強度厚肉化が進んでいる。一般に、鋼板の脆性亀裂伝播停止特性は、高強度あるいは厚肉材になるほど劣化する傾向にあるため、脆性亀裂伝播停止特性への要求も一段と高度化している。 For example, vessels such as container ships and bulk carriers use high-strength thick materials for the hull outer plates due to their structures, but recently, with the increase in size of the hull, the strength has increased further. It is out. Generally, since the brittle crack propagation stop property of a steel sheet tends to deteriorate as the strength or thickness of the steel plate increases, the demand for the brittle crack propagation stop property is further advanced.
ここで、鋼材の脆性亀裂伝播停止特性を向上させる手段として、従来からＮｉ含有量を増加させる方法が知られており、例えば、液化天然ガス（ＬＮＧ）の貯槽タンクにおいては、９％Ｎｉ鋼が商業規模で使用されている。しかしながら、Ｎｉ添加量の増加は、製造コストの大幅な上昇を余儀なくさせるため、ＬＮＧ貯槽タンク以外の用途への適用は難しい。 Here, as a means for improving the brittle crack propagation stopping characteristics of steel materials, a method of increasing the Ni content has been conventionally known. For example, in a storage tank of liquefied natural gas (LNG), 9% Ni steel is used. Used on a commercial scale. However, an increase in the amount of Ni added necessitates a significant increase in manufacturing cost, and therefore it is difficult to apply to applications other than the LNG storage tank.
他方、ＬＮＧのような極低温にまで至らない、例えば、船舶やラインパイプに使用される、板厚が５０ｍｍ未満の比較的薄手の鋼材に対しては、ＴＭＣＰ法により細粒化を図り、低温靭性を向上させることで、優れた脆性亀裂伝播停止特性を実現することができる。 On the other hand, for thin steel materials with a plate thickness of less than 50 mm used for ships and line pipes that do not reach extremely low temperatures, such as LNG, fine graining is attempted by the TMCP method. By improving toughness, excellent brittle crack propagation stopping characteristics can be realized.
また、近年、合金コストを上昇させることなく、鋼板の表層部の組織を超微細化する技術が、脆性亀裂伝播停止特性を向上させる技術として提案されている。 In recent years, a technique for making the microstructure of the surface layer portion of a steel sheet ultrafine without increasing the alloy cost has been proposed as a technique for improving the brittle crack propagation stop characteristic.
たとえば、特許文献１には、脆性亀裂が伝播する際に、鋼材表層部に発生するシアリップ（塑性変形領域）が脆性亀裂伝播停止特性の向上に効果があることに着目し、シアリップ部分の結晶粒を微細化させることで伝播する脆性亀裂が有する伝播エネルギーを吸収させる方法が開示されている。また、特許文献１には、鋼板の製造方法として、熱間圧延後の制御冷却によって表層部分をＡｒ３変態点以下に冷却し、その後制御冷却を停止して表層部分を変態点以上に復熱させる工程を１回以上繰り返して行い、この間に鋼材に圧下を加えることにより、繰り返し変態を生じさせ又は加工再結晶させることで、表層部分に超微細なフェライト組織またはベイナイト組織を生成させる技術が開示されている。For example, Patent Document 1 focuses on the fact that shear lip (plastic deformation region) generated in a steel surface layer portion when brittle crack propagates is effective in improving brittle crack propagation stop characteristics, A method of absorbing propagation energy possessed by a brittle crack propagating by refining the layer is disclosed. In Patent Document 1, as a method of manufacturing a steel sheet, the surface layer portion is cooled to the Ar 3 transformation point or less by controlled cooling after hot rolling, and then the control cooling is stopped to recover the surface layer portion to the transformation point or more. Disclosed is a technique for generating an ultrafine ferrite structure or bainite structure in the surface layer part by repeatedly performing the step of causing the steel material to be subjected to reduction during this time, thereby repeatedly causing transformation or processing recrystallization. Has been.
特許文献２には、フェライト−パーライトを主体のミクロ組織とする鋼材において脆性亀裂伝播停止特性を向上させるために、鋼材の表裏面の表面部を、円相当粒径：５μｍ以下、かつアスペクト比：２以上のフェライト粒を有するフェライト組織を５０％以上有する層で構成し、仕上げ圧延中の１パス当りの最大圧下率を１２％以下とすることで局所的な再結晶現象を抑制し、フェライト粒径のバラツキを抑えることが重要であることが開示されている。 In Patent Document 2, in order to improve brittle crack propagation stopping characteristics in a steel material mainly composed of ferrite-pearlite, the surface portions of the front and back surfaces of the steel material have a circle equivalent particle diameter of 5 μm or less and an aspect ratio: Consists of a layer having 50% or more of a ferrite structure having two or more ferrite grains, and suppressing the local recrystallization phenomenon by setting the maximum rolling reduction per pass during finish rolling to 12% or less. It is disclosed that it is important to suppress variation in diameter.
特許文献３には、フェライト結晶粒の微細化のみならずフェライト結晶粒内に形成されるサブグレインを利用して脆性亀裂伝播停止特性を向上させる、ＴＭＣＰ法を利用した技術が記載されている。具体的には、板厚３０〜４０ｍｍの鋼板を対象とし、鋼板表層の冷却および復熱などの複雑な温度制御を必要とせずに、（ａ）微細なフェライト結晶粒を確保する圧延条件、（ｂ）鋼材板厚の５％以上の部分に微細フェライト組織を生成する圧延条件、（ｃ）微細フェライトに集合組織を発達させるとともに加工（圧延）により導入した転位を熱的エネルギーにより再配置しサブグレインを形成させる圧延条件、および（ｄ）形成した微細なフェライト結晶粒と微細なサブグレイン粒の粗大化を抑制する冷却条件によって脆性亀裂伝播停止特性を向上させることが記載されている。 Patent Document 3 describes a technique using the TMCP method that improves the brittle crack propagation stop characteristics by utilizing subgrains formed in ferrite crystal grains as well as miniaturization of ferrite crystal grains. Specifically, for a steel sheet having a thickness of 30 to 40 mm, without requiring complicated temperature control such as cooling and recuperation of the steel sheet surface layer, (a) rolling conditions for securing fine ferrite crystal grains, ( b) Rolling conditions for generating a fine ferrite structure in a portion of 5% or more of the steel plate thickness, (c) A texture is developed in the fine ferrite, and dislocations introduced by processing (rolling) are rearranged by thermal energy and sub- It is described that brittle crack propagation stopping characteristics are improved by rolling conditions for forming grains and (d) cooling conditions for suppressing coarsening of the formed fine ferrite crystal grains and fine subgrain grains.
また、制御圧延において、変態したフェライトに圧下を加えて集合組織を発達させることにより、脆性亀裂伝播停止特性を向上させる方法も知られている。鋼材の破壊面上にセパレーションを板厚方向と平行な方向に生ぜしめ、脆性亀裂先端の応力を緩和させることにより、脆性破壊に対する抵抗を高めるものである。 In addition, in controlled rolling, a method of improving the brittle crack propagation stop property by applying a reduction to a transformed ferrite to develop a texture is also known. Separation occurs on the fracture surface of the steel material in a direction parallel to the plate thickness direction to relieve stress at the brittle crack tip, thereby increasing resistance to brittle fracture.
例えば、特許文献４には、制御圧延により（１１０）面Ｘ線強度比を２以上とし、かつ円相当径２０μｍ以上の粗大粒の面積率を１０％以下とすることにより、耐脆性破壊特性を向上させた鋼板が記載されている。 For example, Patent Document 4 discloses that the (110) plane X-ray intensity ratio is 2 or more by controlled rolling, and the area ratio of coarse grains having a circle-equivalent diameter of 20 μm or more is 10% or less. An improved steel sheet is described.
特許文献５には、継手部の脆性亀裂伝播停止性能の優れた溶接構造用鋼として、集合組織発達により応力負荷方向と亀裂伝播方向をずらすため、板厚内部の圧延面での（１００）面のＸ線面強度比を１．５以上とした鋼板が開示されている。 In Patent Document 5, as a steel for welded structure having excellent brittle crack propagation stopping performance of a joint part, the stress load direction and the crack propagation direction are shifted by texture development. A steel sheet having an X-ray surface strength ratio of 1.5 or more is disclosed.
特許文献６には、板厚の５０％以上の領域において板面と平行な面での（２１１）面のＸ線回折強度比を１．５以上とすることにより、靭性および脆性亀裂伝播停止特性に優れた薄肉厚鋼板が記載されている。 Patent Document 6 discloses toughness and brittle crack propagation stop characteristics by setting the X-ray diffraction intensity ratio of (211) plane in a plane parallel to the plate surface to 1.5 or more in a region of 50% or more of the plate thickness. Describes a thin-walled thick steel plate.
特許文献７には、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とについて、それぞれ、圧延面と平行な（１００）Ｘ線面強度比、及び、圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比を規定した集合組織を有する脆性亀裂伝播停止性能に優れた高強度厚鋼板が記載されている。 In Patent Document 7, (100) X-ray plane intensity ratio parallel to the rolling surface, respectively, for the front and back layer portions up to 25% of the plate thickness from the front and back surfaces of the steel plate and the other plate thickness center portion, and In addition, a high-strength thick steel plate excellent in brittle crack propagation stopping performance having a texture defining a (111) and / or (211) X-ray plane strength ratio parallel to the rolling surface is described.
特許文献８には、鋼板の板厚方向の中心を中央として板厚の１０％以上、５０％未満の中心部領域において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比を規定し、さらに、前記中心部領域より表裏面側の表裏面領域において圧延面と平行な（１１１）又は／及び（２１１）Ｘ線面強度比を規定した集合組織を有する脆性き裂伝播停止性能に優れた高強度厚鋼板が記載されている。 In Patent Document 8, (111) or / and (211) X-ray plane intensity ratio parallel to the rolling surface in the central region of 10% or more and less than 50% of the plate thickness with the center in the plate thickness direction of the steel plate as the center. Furthermore, in the front and back surface regions on the front and back sides from the central region, the brittle crack propagation stop having a texture that defines the (111) or / and (211) X-ray surface strength ratio parallel to the rolling surface A high strength thick steel plate with excellent performance is described.
特許文献９〜１２には、板厚中央部および板厚１／４部における各種Ｘ線面強度比を規定した集合組織を有する構造用高強度厚鋼板が記載されている。 Patent Documents 9 to 12 describe structural high-strength thick steel plates having a texture that defines various X-ray plane strength ratios in the center portion of the plate thickness and the ¼ portion of the plate thickness.
上述したように、脆性亀裂伝播停止性能に優れた鋼板やその製造方法に関して種々の提案がなされているが、大型構造物に使用される鋼材には安全性の観点から、優れた溶接熱影響部の靭性、特にボンド部の靭性に優れることも同時に要求される。 As described above, various proposals have been made regarding steel sheets excellent in brittle crack propagation stopping performance and methods for producing the same, but from the viewpoint of safety, steel materials used in large structures have excellent weld heat affected zone. It is also required to have excellent toughness, particularly toughness at the bond part.
ボンド部は、大入熱溶接時の融点直下の高温にさらされて、オーステナイト結晶粒が最も粗大化しやすく、その後の冷却によって脆弱な上部ベイナイト組織に変態し、更に、ウィドマンステッテン組織や島状マルテンサイトが生成して靭性が低下する。 The bond part is exposed to a high temperature just below the melting point during high heat input welding, the austenite crystal grains are most likely to be coarsened, and then transformed into a fragile upper bainite structure by cooling. Martensite is formed and toughness is reduced.
ボンド部の靭性向上に関しては種々の研究がなされ、例えば、ＴｉＮの微細分散によるオーステナイトの粗大化抑制やフェライト変態核としての利用のほか、希土類元素（ＲＥＭ）をＴｉと複合添加することにより、鋼中に微細粒子を分散させてオーステナイトの粒成長を防止し、溶接部の靭性向上を図る方法が提案されている（特許文献１３、１４）。 Various studies have been made on improving the toughness of the bond part. For example, by adding a rare earth element (REM) with Ti in addition to suppressing the austenite coarsening by fine dispersion of TiN and using it as a ferrite transformation nucleus, A method has been proposed in which fine particles are dispersed therein to prevent austenite grain growth and to improve the toughness of the welded portion (Patent Documents 13 and 14).
また、Ｔｉ酸化物やＭｇ酸化物を利用したり（特許文献１５、１６）、ＢＮによりフェライト核を生成したり、ＣａやＲＥＭを添加することで硫化物の形態を制御して、靭性を向上させることが提案されている。その他Ｃａ、Ｏ、Ｓ量を制御し、ＣａおよびＭｎの複合硫化物をフェライト核とし微細に分散させることによって、靭性を向上させる方法が提案されている（特許文献１７）。 In addition, Ti oxide and Mg oxide are used (Patent Documents 15 and 16), ferrite nuclei are generated by BN, and the form of sulfide is controlled by adding Ca or REM to improve toughness. It has been proposed to let In addition, a method for improving toughness by controlling the amounts of Ca, O, and S and finely dispersing Ca and Mn composite sulfides as ferrite nuclei has been proposed (Patent Document 17).
また、特許文献１８には集合組織を制御することで脆性亀裂伝播停止特性を向上させた大入熱溶接用鋼およびその製造方法が開示されている。 Patent Document 18 discloses a steel for high heat input welding in which the brittle crack propagation stopping characteristics are improved by controlling the texture, and a manufacturing method thereof.
しかし、特許文献１、２に記載の脆性亀裂伝播停止特性に優れた鋼材は、鋼材表層部のみを一旦冷却した後に復熱させ、かつ復熱中に加工を加えることによって、特定の組織を得るもので、実生産規模では制御が容易でなく、圧延、冷却設備への負荷が大きいプロセスである。 However, the steel materials excellent in brittle crack propagation stopping characteristics described in Patent Documents 1 and 2 are obtained by recooling only the steel surface layer part and then recovering the heat, and by applying processing during the recuperation, a specific structure is obtained. In actual production scale, it is a process that is difficult to control and has a heavy load on rolling and cooling equipment.
また、非特許文献１では、板厚：６５ｍｍの鋼板の脆性亀裂伝播停止特性が評価され、母材の大型脆性亀裂伝播停止試験で脆性亀裂が停止しない結果が報告されている。さらに、非特許文献１では、供試材のＥＳＳＯ試験において使用温度−１０℃におけるＫｃａの値が３０００Ｎ／ｍｍ３／２に満たない結果を示し、５０ｍｍを超える板厚の鋼板を適用した船体構造の場合、安全性確保が課題となることが示唆されている。Further, in Non-Patent Document 1, the brittle crack propagation stop property of a steel plate having a plate thickness of 65 mm is evaluated, and a result that a brittle crack does not stop in a large brittle crack propagation stop test of a base material is reported. Furthermore, Non-Patent Document 1 shows a result that the value of Kca at an operating temperature of −10 ° C. is less than 3000 N / mm 3/2 in the ESSO test of the specimen, and a hull structure to which a steel plate having a thickness exceeding 50 mm is applied. In this case, it is suggested that ensuring safety is an issue.
これに対して、上述した特許文献１〜６に記載された鋼板は、いずれも、製造条件や開示されている実験データから、板厚５０ｍｍ程度が主な対象であって、７０ｍｍもしくはそれ以上の厚肉材への適用については、所定の特性が得られるかが不明で、船体構造において必要な、板厚方向の亀裂伝播特性に対しては不明である。 On the other hand, the steel sheets described in Patent Documents 1 to 6 described above are mainly subject to a plate thickness of about 50 mm from the manufacturing conditions and the disclosed experimental data, and are 70 mm or more. Regarding application to thick materials, it is unclear whether a predetermined characteristic can be obtained, and it is unclear as to the crack propagation characteristic in the thickness direction, which is necessary in the hull structure.
特許文献７および８には、温度勾配型の標準ＥＳＳＯ試験によりアレスト性が評価されているが、溶接部靭性の評価はなされておらず、入熱量が３００ｋＪ／ｃｍを超える大入熱溶接の適用可否についても不明である。 In Patent Documents 7 and 8, arrestability is evaluated by a temperature gradient type standard ESSO test, but the evaluation of weld toughness has not been made, and the application of large heat input welding in which the heat input exceeds 300 kJ / cm Whether it is possible is unknown.
特許文献９〜１２に開示された技術は、Ａｒ３点未満の温度域、すなわち、フェライト−オーステナイト二相域での圧延を必須としている。このため、高精度の圧延技術を必要とするだけでなく、通常よりも低温域での圧延であるため、生産能率が低下し、また、鋼板形状を平坦にするにも特段の配慮が必要とされる。このため、生産性を犠牲にしない製造条件にて、優れた脆性亀裂伝播停止特性を確保する技術が望まれる。The techniques disclosed in Patent Documents 9 to 12 require rolling in a temperature range lower than the Ar 3 point, that is, a ferrite-austenite two-phase range. For this reason, not only high-precision rolling technology is required, but also rolling in a lower temperature range than usual, production efficiency is reduced, and special consideration is also required for flattening the steel plate shape. Is done. For this reason, a technique for securing excellent brittle crack propagation stopping characteristics under manufacturing conditions that do not sacrifice productivity is desired.
一方、溶接施工において、板厚５０ｍｍ以上の厚鋼板を溶接する場合、入熱３００ｋＪ／ｃｍを超える大入熱溶接の適用が検討され、さらなる大入熱化が予想される。 On the other hand, when welding thick steel plates having a thickness of 50 mm or more in welding construction, application of high heat input welding with a heat input exceeding 300 kJ / cm is studied, and further increase in heat input is expected.
しかしながら、特許文献１３、１４記載の、ＴｉＮを主体に利用する技術においてはＴｉＮが溶解する温度域に加熱される溶接部でその作用が消失し、また固溶ＴｉおよびＮにより組織が脆化して著しく靭性が低下するので、３００ｋＪ／ｃｍを超える大入熱溶接部では十分な靭性が得られないことが予想される。 However, in the techniques mainly using TiN described in Patent Documents 13 and 14, the action disappears in the weld zone heated to a temperature range where TiN dissolves, and the structure is embrittled by solute Ti and N. Since the toughness is remarkably lowered, it is expected that sufficient toughness cannot be obtained at a high heat input weld portion exceeding 300 kJ / cm.
さらに、特許文献１５、１６記載の技術のように、Ｔｉ酸化物やＭｇ酸化物を利用してＨＡＺ靭性を改善する場合、これらの酸化物を十分均質に微細分散することは容易でなく、またＣａやＲＥＭを添加する技術においても３００ｋＪ／ｃｍを超える大入熱溶接では溶接熱影響部の高靭性を確保することは困難であった。 Furthermore, as in the techniques described in Patent Documents 15 and 16, when using a Ti oxide or Mg oxide to improve HAZ toughness, it is not easy to finely and uniformly disperse these oxides. Even in the technique of adding Ca or REM, it has been difficult to ensure high toughness of the heat affected zone by high heat input welding exceeding 300 kJ / cm.
また、特許文献１７においては、ＣａおよびＭｎの複合硫化物を利用することで４００ｋＪ／ｃｍを超える溶接熱影響部の靭性を確保しているが、脆性亀裂伝播停止性能に関する検討はなされていない。 In Patent Document 17, the toughness of the weld heat-affected zone exceeding 400 kJ / cm is ensured by using a composite sulfide of Ca and Mn, but no investigation has been made regarding brittle crack propagation stopping performance.
これに対し、特許文献１８においては、溶接熱影響部の靱性を確保するとともに脆性亀裂伝播停止性能に優れる鋼板が開示されている。しかし、本技術を用いたとしても７０ｍｍを超える板厚においては大入熱溶接部の靱性および脆性亀裂伝播特性十分でない場合があった。 On the other hand, Patent Document 18 discloses a steel sheet that ensures the toughness of the weld heat affected zone and has excellent brittle crack propagation stopping performance. However, even when this technology is used, the toughness and brittle crack propagation characteristics of large heat input welds may not be sufficient when the plate thickness exceeds 70 mm.
本発明は、大入熱溶接部の靭性および脆性亀裂伝播停止特性に優れた高強度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。 An object of this invention is to provide the high strength thick steel plate excellent in the toughness of a high heat-input welding part, and the brittle crack propagation stop characteristic, and its manufacturing method.
本発明者らは、上記課題の達成のため鋭意研究を重ね、厚鋼板においても優れた脆性亀裂伝播停止特性を有し、かつ大入熱溶接部の靭性に優れた高強度厚鋼板および当該鋼板を安定して得る製造方法について以下の知見を得た。 The inventors of the present invention have made extensive studies to achieve the above-mentioned problems, and have high brittle crack propagation stopping characteristics even in thick steel sheets, and high strength thick steel sheets excellent in toughness of large heat input welds and the steel sheets. The following knowledge was obtained about the manufacturing method which obtains stably.
（Ｉ）．板厚５０ｍｍを超える厚鋼板において脆性亀裂伝播停止特性に及ぼす集合組織の影響を詳細に調べた結果、鋼板表面の板面（圧延面）に平行な面における（２１１）面Ｘ線強度比が１．２以上、板厚中央（板厚１／２位置）の板面（圧延面）に平行な面における（２１１）面Ｘ線強度比が１．５以上かつ（２２２）面Ｘ線強度比が２．５以下となるように制御することにより優れた脆性亀裂伝播停止特性が得られる。 (I). As a result of investigating the effect of texture on the brittle crack propagation stopping characteristics in a thick steel plate exceeding 50 mm thick, the (211) plane X-ray intensity ratio in a plane parallel to the plate surface (rolled surface) of the steel plate surface is 1. .2 or more, the (211) plane X-ray intensity ratio is 1.5 or more and the (222) plane X-ray intensity ratio in a plane parallel to the plate surface (rolled surface) at the plate thickness center (plate thickness 1/2 position). By controlling to be 2.5 or less, excellent brittle crack propagation stopping characteristics can be obtained.
（ＩＩ）．上記の集合組織は、特定の化学成分組成の鋼をオーステナイト域で圧延を完了する場合において、特定の熱間圧延条件に従うことで得ることができる。上記オーステナイト域での圧延においては、圧延時の温度が低温であるほど高い靭性値と発達した（２１１）面集合組織が得られる。しかしながら、板厚が５０ｍｍを超えるような場合には、圧延時の板厚中央と鋼板表面での温度差が大きくなるため、温度を下げすぎると表層部にフェライト組織が生成してしまい靭性を劣化させる問題が発生する。一方で表層の靭性劣化を抑えるためには圧延温度を上げる必要があるが、その場合には板厚中央での靭性／およびまたは集合組織の発達が不十分となる場合があった。そこで、圧延途中で鋼素材の表裏面を加熱することで、板厚方向の温度分布を制御し、上記の問題を解決することで従来よりも優れた脆性亀裂伝播停止特性を得ることができる。 (II). The above texture can be obtained by following specific hot rolling conditions when rolling a steel having a specific chemical composition in the austenite region. In rolling in the austenite region, the lower the temperature during rolling, the higher the toughness value and the developed (211) plane texture. However, when the plate thickness exceeds 50 mm, the temperature difference between the center of the plate thickness during rolling and the surface of the steel plate becomes large, so if the temperature is lowered too much, a ferrite structure is formed in the surface layer and deteriorates toughness. Cause problems. On the other hand, in order to suppress deterioration of the toughness of the surface layer, it is necessary to raise the rolling temperature. In that case, the development of toughness and / or texture at the center of the sheet thickness may be insufficient. Therefore, by heating the front and back surfaces of the steel material in the middle of rolling, the temperature distribution in the plate thickness direction is controlled and the above-described problems can be solved to obtain a brittle crack propagation stopping characteristic superior to that of the prior art.
（ＩＩＩ）．上記特定の化学成分組成の鋼板の溶接ボンド部の靭性は脆化組織に影響され、この脆化組織の靭性は溶接後の冷却時にフェライト変態を促進させる変態核の微細化を行う事で大きく向上する。変態核を微細に分散させるためには、添加量を下記の（１）式を満足するようにＣａ、Ｓ、Ｏ量を調節する。
０＜［（Ｃａ−（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５］／Ｓ＜１ ・・・（１）
すなわち、鋼を溶製する際の凝固段階でＣａＳを晶出させるにあたり、（１）式を満足するようにＣａ、Ｓの添加量および添加時の溶鋼中の溶存酸素量を制御することによって、ＣａＳの晶出後の固溶Ｓ量を確保すれば、ＣａＳの表面上にＭｎＳが析出する。ＭｎＳはフェライト核生成能を有し、その周囲にＭｎの希薄帯が形成されるとフェライト変態が促進され、溶接熱影響部の靭性を向上させる。ＭｎＳ上にＴｉＮ、ＢＮ、ＡｌＮ等のフェライト生成核が析出することによって、より一層、フェライト変態が促進される。(III). The toughness of the weld bond part of the steel sheet with the above-mentioned specific chemical composition is affected by the embrittlement structure, and the toughness of this embrittlement structure is greatly improved by refining the transformation nucleus that promotes the ferrite transformation during cooling after welding. To do. In order to finely disperse the transformation nuclei, the amounts of Ca, S, and O are adjusted so as to satisfy the following formula (1).
0 <[(Ca− (0.18 + 130 × Ca) × O) /1.25] / S <1 (1)
That is, in crystallization of CaS in the solidification stage when melting steel, by controlling the amount of Ca and S added and the amount of dissolved oxygen in the molten steel at the time of addition so as to satisfy the formula (1), If the amount of solid solution S after crystallization of CaS is secured, MnS will precipitate on the surface of CaS. MnS has a ferrite nucleation ability, and when a Mn dilute band is formed around it, ferrite transformation is promoted and the toughness of the weld heat affected zone is improved. Ferrite transformation is further promoted by precipitation of ferrite-forming nuclei such as TiN, BN, and AlN on MnS.
以上、前述した化学成分および製造プロセスを用い、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比が１．２以上、板厚中央（板厚１／２位置）の板面に平行な面における（２１１）面Ｘ線強度比が１．５以上かつ（２２２）面Ｘ線強度比が２．５以下の集合組織を有する場合に、極めて優れた脆性亀裂伝播停止特性が得られ、かつ大入熱溶接におけるボンド部の靱性に優れることを知見した。 As described above, the (211) plane X-ray intensity ratio in the plane parallel to the plate surface of the steel plate surface is 1.2 or more and the plate surface at the plate thickness center (plate thickness 1/2 position) using the chemical composition and the manufacturing process described above. When the (211) plane X-ray intensity ratio is 1.5 or more and the (222) plane X-ray intensity ratio is 2.5 or less in a plane parallel to the surface, extremely excellent brittle crack propagation stop characteristics are obtained. In addition, it has been found that the bond portion has high toughness in high heat input welding.
本発明は、上記した知見に、さらに検討を加えて完成されたものであって、本発明の要旨構成は次のとおりである。
［１］成分組成が、質量％で、Ｃ：０．０３〜０．２０％、Ｓｉ：０．０１〜０．３０％、Ｍｎ：１．５〜３．０％、Ｐ：０．０２％以下、Ｓ：０．０００５〜０．０１％、Ｔｉ：０．００５〜０．０３０％、Ａｌ：０．００５〜０．０８０％、Ｎ：０．００２５〜０．００７５％、Ｃａ：０．０００３〜０．００３０％、Ｂ：０．０００３〜０．００３０％、Ｏ：０．００３０％以下を含有し、かつ、Ｃａ、Ｏ、Ｓが下記（１）式を満たし、下記（２）式で定義されるＣｅｑが０．３６〜０．５０の範囲にあり、残部がＦｅおよび不可避的不純物からなり、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比が１．２以上、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比が１．５以上かつ（２２２）面Ｘ線強度比が２．５以下であり、
３００ｋＪ／ｃｍ超の入熱量で溶接した溶接継手ボンド部の−４０℃におけるシャルピー吸収エネルギー（ｖＥ−４０）が８０Ｊ以上であり、ＥＳＳＯ試験による−１０℃におけるＫｃａ値（Ｋｃａ（−１０℃））が６０００Ｎ／ｍｍ３／２以上であり、引張り強さが５８０ＭＰａ以上であり、板厚が５０ｍｍ超である、高強度厚鋼板。
０＜［（Ｃａ−（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５］／Ｓ＜１ ・・・（１）
Ｃｅｑ＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５ ・・・（２）
ただし、上記（１）式および（２）式における各元素記号は各元素の含有量（質量％）を表し、含有しない元素は０とする。
［２］シャルピー衝撃試験による板厚中央部の破面遷移温度（ｖＴｒｓ）が−６０℃以下である、［１］に記載の高強度厚鋼板。
［３］シャルピー衝撃試験による鋼板表面から板厚方向に５ｍｍの位置の破面遷移温度（ｖＴｒｓ）が−６０℃以下である、［１］または［２］に記載の高強度厚鋼板。
［４］成分組成が、さらに、質量％で、Ｎｂ：０．００３〜０．０４０％、Ｃｕ：０．０１〜０．５％、Ｎｉ：０．０１〜２．０％、Ｃｒ：０．０１〜０．５％、Ｍｏ：０．０１〜０．５％のなかから選ばれる１種または２種以上を含有する、［１］〜［３］のいずれかに記載の高強度厚鋼板。
［５］成分組成が、さらに、質量％で、Ｖ：０．００１〜０．１０％、Ｍｇ：０．０００５〜０．００５０％、Ｚｒ：０．０００５〜０．０２００％、ＲＥＭ：０．０００５〜０．０２００％のなかから選ばれる１種または２種以上を含有する、［１］〜［４］のいずれかに記載の高強度厚鋼板。
［６］前記［１］〜［５］のいずれかに記載の高強度厚鋼板の製造方法であって、前記成分組成を有する鋼素材を、１０００〜１２００℃の温度に加熱したのち、板厚中央が、オーステナイト再結晶温度域、次いで、オーステナイト未再結晶温度域で熱間圧延を行い、少なくとも前記オーステナイト未再結晶温度域での熱間圧延の一部を、鋼板表面の温度−板厚中央の温度≧−４０℃となる条件で行うように板厚方向の温度分布を制御し、熱間圧延終了時における鋼板表面の温度と板厚中央の温度との温度差を５℃以内とし、かつ、板厚中央の温度をＡｒ３℃〜（Ａｒ３+３０）℃とする、高強度厚鋼板の製造方法。The present invention has been completed by further studying the above knowledge, and the gist of the present invention is as follows.
[1] Component composition is mass%, C: 0.03 to 0.20%, Si: 0.01 to 0.30%, Mn: 1.5 to 3.0%, P: 0.02% Hereinafter, S: 0.0005 to 0.01%, Ti: 0.005 to 0.030%, Al: 0.005 to 0.080%, N: 0.0025 to 0.0075%, Ca: 0.00. 0003 to 0.0030%, B: 0.0003 to 0.0030%, O: 0.0030% or less, and Ca, O, and S satisfy the following formula (1), and the following formula (2) The Ceq defined by the above is in the range of 0.36 to 0.50, the balance is made of Fe and inevitable impurities, and the (211) plane X-ray intensity ratio in the plane parallel to the plate surface of the steel plate surface is 1.2. As described above, the (211) plane X-ray intensity ratio is 1.5 or more and the (222) plane X-ray intensity ratio is 2 in a plane parallel to the plate surface at the center of the sheet thickness. 5 or less,
The Charpy absorbed energy (vE -40 ) at -40 ° C of the welded joint part welded with a heat input of more than 300 kJ / cm is 80 J or more, and the Kca value (Kca (-10 ° C)) at -10 ° C by the ESSO test. There is a 6000 N / mm 3/2 or more and tensile strength of 580MPa or more, the thickness is 50mm greater, high strength thick steel plate.
0 <[(Ca− (0.18 + 130 × Ca) × O) /1.25] / S <1 (1)
Ceq = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (2)
However, each element symbol in the above formulas (1) and (2) represents the content (mass%) of each element, and the elements not contained are 0.
[2] The high-strength thick steel plate according to [1], wherein a fracture surface transition temperature (vTrs) at a central portion of the plate thickness by a Charpy impact test is −60 ° C. or lower.
[3] The high-strength thick steel plate according to [1] or [2], wherein a fracture surface transition temperature (vTrs) at a position 5 mm from the steel plate surface in the thickness direction by a Charpy impact test is −60 ° C. or lower.
[4] The component composition is further mass%, Nb: 0.003 to 0.040%, Cu: 0.01 to 0.5%, Ni: 0.01 to 2.0%, Cr: 0.00. The high-strength thick steel plate according to any one of [1] to [3], containing one or more selected from 01 to 0.5% and Mo: 0.01 to 0.5%.
[5] The component composition is further in mass%, V: 0.001 to 0.10%, Mg: 0.0005 to 0.0050%, Zr: 0.0005 to 0.0200%, REM: 0.00. The high-strength thick steel plate according to any one of [1] to [4], containing one or more selected from 0005 to 0.0200%.
[6] The method for producing a high-strength thick steel plate according to any one of [1] to [5], wherein the steel material having the above composition is heated to a temperature of 1000 to 1200 ° C. The center is hot-rolled in the austenite recrystallization temperature range, then in the austenite non-recrystallization temperature range, and at least part of the hot rolling in the austenite non-recrystallization temperature range is performed at the temperature-sheet thickness center of the steel sheet surface. The temperature distribution in the sheet thickness direction is controlled so as to be performed under the condition of temperature ≧ −40 ° C., and the temperature difference between the surface temperature of the steel sheet and the temperature at the center of the sheet thickness at the end of hot rolling is within 5 ° C., and A method for producing a high-strength thick steel plate, wherein the temperature at the center of the plate thickness is Ar 3 ° C. to (Ar 3 +30) ° C.
本発明によれば、大入熱溶接部の靭性および脆性亀裂伝播停止特性に優れた高強度厚鋼板およびその製造方法を提供することができる。 ADVANTAGE OF THE INVENTION According to this invention, the high strength thick steel plate excellent in the toughness of a high heat-input welding part and a brittle crack propagation stop characteristic, and its manufacturing method can be provided.
本発明によれば、板厚が５０ｍｍを超える厚鋼板においても、圧延条件を最適化し、板厚方向での集合組織を制御することができる工業的に簡易なプロセスにより、安定して入熱量が３００ｋＪ／ｃｍ超の大入熱溶接部の靭性および脆性亀裂伝播停止特性に優れた高強度厚鋼板を提供することができる。 According to the present invention, even in a thick steel plate having a plate thickness exceeding 50 mm, the heat input can be stably stabilized by an industrially simple process that can optimize the rolling conditions and control the texture in the plate thickness direction. It is possible to provide a high-strength thick steel plate excellent in toughness and brittle crack propagation stopping characteristics of a high heat input weld zone exceeding 300 kJ / cm.
本発明によれば、板厚方向の各位置に応じて集合組織および靭性値が適切に制御されるので、脆性亀裂伝播停止特性に優れ、例えば、造船分野では、コンテナ船、バルクキャリアーの強力甲板部構造においてハッチサイドコーミングに接合される甲板部材へ適用することにより、船舶の安全性向上に寄与するところが大で産業上極めて有用である。 According to the present invention, the texture and toughness value are appropriately controlled according to the respective positions in the plate thickness direction, so that it has excellent brittle crack propagation stopping characteristics. For example, in the shipbuilding field, a strong deck of a container ship and a bulk carrier. The application to the deck member joined to the hatch side combing in the part structure is very useful industrially because it contributes to improving the safety of the ship.
以下、本発明を具体的に説明する。
本発明では、成分組成および鋼板内部の集合組織を規定する。Hereinafter, the present invention will be specifically described.
In the present invention, the component composition and the texture inside the steel sheet are defined.
［成分組成］
以下の説明において、鋼板成分における％は、すべて質量％を意味する。[Ingredient composition]
In the following description, “%” in the steel plate component means “% by mass”.
Ｃ：０．０３〜０．２０％
Ｃは、鋼の強度を向上する元素であり、本発明では、所望の強度を確保するために、Ｃの含有量を０．０３％以上とする。一方、Ｃの含有量が０．２０％を超えると、溶接性が劣化するばかりか靭性にも悪影響がある。このため、Ｃの含有量は、０．０３〜０．２０％の範囲に規定する。なお、Ｃの含有量は、好ましくは０．０４％以上であり、より好ましくは０．０５％以上である。また、Ｃの含有量は、好ましくは０．１５％以下であり、より好ましくは０．１２％以下である。C: 0.03-0.20%
C is an element that improves the strength of steel, and in the present invention, the C content is set to 0.03% or more in order to ensure a desired strength. On the other hand, if the C content exceeds 0.20%, not only the weldability is deteriorated but also the toughness is adversely affected. For this reason, content of C is prescribed | regulated in the range of 0.03-0.20%. The C content is preferably 0.04% or more, and more preferably 0.05% or more. Further, the C content is preferably 0.15% or less, more preferably 0.12% or less.
Ｓｉ：０．０１〜０．３０％
Ｓｉは、脱酸元素として、また、鋼の強化元素として有効であるが、０．０１％未満の含有量ではその効果が十分に得られない。一方、０．３０％を超えると鋼板の表面性状を損なうばかりか、母材および大入熱溶接継手のボンド部の靭性が極端に劣化する。従って、その含有量を０．０１〜０．３０％の範囲とする。Ｓｉの含有量の上限値としては、０．２０％が好ましく、０．１０％がより好ましい。Si: 0.01-0.30%
Si is effective as a deoxidizing element and as a steel strengthening element, but if the content is less than 0.01%, the effect cannot be sufficiently obtained. On the other hand, if it exceeds 0.30%, not only the surface properties of the steel sheet are impaired, but also the toughness of the bond portion of the base material and the high heat input welded joint is extremely deteriorated. Therefore, the content is made 0.01 to 0.30%. The upper limit of the Si content is preferably 0.20%, and more preferably 0.10%.
Ｍｎ：１．５〜３．０％
Ｍｎは、強化元素および焼入れ元素として含有する。Ｍｎの含有量が、１．５％より少ないとその効果が十分でない一方で、３．０％を超えると溶接性が劣化し、鋼材コストも上昇する。そのため、Ｍｎの含有量は、１．５〜３．０％とする。Ｍｎの含有量は、好ましくは１．７％以上であり、より好ましくは１．９％以上である。また、Ｍｎの含有量は、好ましくは２．７％以下であり、より好ましくは２．５％以下である。Mn: 1.5 to 3.0%
Mn is contained as a strengthening element and a quenching element. If the Mn content is less than 1.5%, the effect is not sufficient. On the other hand, if it exceeds 3.0%, the weldability deteriorates and the steel material cost also increases. Therefore, the Mn content is 1.5 to 3.0%. The Mn content is preferably 1.7% or more, and more preferably 1.9% or more. Further, the Mn content is preferably 2.7% or less, and more preferably 2.5% or less.
Ｐ：０．０２％以下
Ｐは、含有量が多くなると靭性が劣化してしまう。板厚：５０ｍｍ超の鋼板に対して、良好な靭性を保つためには、Ｐの含有量を０．０２％以下に制御する。好ましくは、Ｐの含有量を０．０１％以下、さらに好ましくは０．００６％以下に制御する。P: 0.02% or less P increases the toughness as the content increases. Thickness: In order to maintain good toughness with respect to steel sheets exceeding 50 mm, the P content is controlled to 0.02% or less. Preferably, the P content is controlled to 0.01% or less, more preferably 0.006% or less.
Ｓ：０．０００５〜０．０１％
Ｓは、含有量が多くなると靭性が劣化してしまう。そのため、Ｓは０．０１％以下に抑制する。一方、大入熱溶接継手のボンド部おいて優れた靱性を得るためには、Ｓを０．０００５％以上含有することが必要である。Ｓの含有量は、好ましくは０．００１０％以上である。また、Ｓの含有量は、好ましくは０．００５０％以下であり、より好ましくは０．００３０％以下である。S: 0.0005 to 0.01%
If S content increases, toughness will deteriorate. Therefore, S is suppressed to 0.01% or less. On the other hand, in order to obtain excellent toughness at the bond part of the high heat input welded joint, it is necessary to contain 0.0005% or more of S. The content of S is preferably 0.0010% or more. Further, the content of S is preferably 0.0050% or less, and more preferably 0.0030% or less.
Ｔｉ：０．００５〜０．０３０％
Ｔｉは、微量の含有により、窒化物、炭化物、あるいは炭窒化物を形成し、結晶粒を微細化して母材靭性を向上させる効果を有する。また、Ｔｉは、凝固時にＴｉＮとなって析出し、溶接部でのオーステナイトの粗大化抑制やフェライト変態核となって高靭性化に寄与する。Ｔｉの含有量が、０．００５％未満ではその効果が少なく、一方、０．０３０％を超えるとＴｉＮ粒子の粗大化によってその効果が得られなくなるため、Ｔｉの含有量は０．００５〜０．０３０％とする。Ｔｉの含有量は、好ましくは０．００８％以上である。また、Ｔｉの含有量は、好ましくは０．０２０％以下である。Ti: 0.005-0.030%
Ti has the effect of forming nitrides, carbides, or carbonitrides by adding a trace amount, and making the crystal grains finer to improve the base material toughness. Further, Ti precipitates as TiN during solidification and contributes to high toughness by suppressing the austenite coarsening in the weld and ferrite transformation nuclei. If the Ti content is less than 0.005%, the effect is small. On the other hand, if the Ti content exceeds 0.030%, the effect cannot be obtained due to the coarsening of TiN particles, so the Ti content is 0.005 to 0. 0.030%. The Ti content is preferably 0.008% or more. Further, the Ti content is preferably 0.020% or less.
Ａｌ：０．００５〜０．０８０％
Ａｌは、脱酸剤として作用し、このためには０．００５％以上の含有を必要とするが、０．０８０％を超えて含有すると、靭性を低下させるとともに、溶接した場合に、溶接金属部の靭性を低下させる。このため、Ａｌの含有量は、０．００５〜０．０８０％の範囲に規定した。なお、Ａｌの含有量は、好ましくは０．０２％以上である。また、Ａｌの含有量は、好ましくは０．０６０％以下である。Al: 0.005-0.080%
Al acts as a deoxidizer, and for this purpose, it needs to contain 0.005% or more. However, if it contains more than 0.080%, the toughness is lowered and, when welded, weld metal Reduce the toughness of the part. For this reason, the content of Al is specified in the range of 0.005 to 0.080%. The Al content is preferably 0.02% or more. Further, the Al content is preferably 0.060% or less.
Ｎ：０．００２５〜０．００７５％
Ｎは、ＴｉＮの必要量を確保するために必要な元素で、０．００２５％未満では十分なＴｉＮ量が得られず、０．００７５％を超えると溶接熱サイクルによってＴｉＮが溶解する領域において固溶Ｎ量が増加して靭性を著しく低下させるため、０．００２５〜０．００７５％とする。Ｎの含有量は、好ましくは０．００３０％以上である。また、Ｎの含有量は、好ましくは０．００７０％以下である。N: 0.0025 to 0.0075%
N is an element necessary for securing the necessary amount of TiN. If it is less than 0.0025%, a sufficient amount of TiN cannot be obtained, and if it exceeds 0.0075%, it is solidified in the region where TiN is dissolved by the welding heat cycle. In order to increase the amount of dissolved N and significantly reduce toughness, the content is made 0.0025 to 0.0075%. The N content is preferably 0.0030% or more. The N content is preferably 0.0070% or less.
Ｃａ：０．０００３％〜０．００３０％
Ｃａは、Ｓの固定による靭性改善効果を有する元素である。このような効果を発揮させるためには少なくとも０．０００３％は含有することが必要であるが、０．００３０％を超えて含有しても効果が飽和するため、０．０００３〜０．００３０％とする。Ｃａの含有量は、好ましくは０．００１０％以上である。また、Ｃａの含有量は、好ましくは０．００２５％以下である。Ca: 0.0003% to 0.0030%
Ca is an element having an effect of improving toughness by fixing S. In order to exert such an effect, it is necessary to contain at least 0.0003%, but even if it contains more than 0.0030%, the effect is saturated, so 0.0003 to 0.0030% And The Ca content is preferably 0.0010% or more. Further, the Ca content is preferably 0.0025% or less.
Ｂ：０．０００３〜０．００３０％
Ｂは、溶接熱影響部でＴｉＮの溶解によるＮをＢＮとして固定し、溶接部靭性の劣化を抑制する。また、焼入性を向上させ母材の強度確保に有効に寄与する。このような効果は０．０００３％以上の添加で発揮され、また、０．００３０％よりも多く添加してもその効果は飽和するため、Ｂの含有量は０．０００３〜０．００３０％とする。Ｂの含有量は、好ましくは０．０００８％以上である。また、Ｂの含有量は、好ましくは０．００２５％以下である。B: 0.0003 to 0.0030%
B fixes N due to dissolution of TiN as BN in the weld heat affected zone, and suppresses the deterioration of the weld toughness. Moreover, it improves hardenability and contributes effectively to securing the strength of the base material. Such an effect is exhibited by addition of 0.0003% or more, and even if added more than 0.0030%, the effect is saturated. Therefore, the content of B is 0.0003 to 0.0030%. To do. The content of B is preferably 0.0008% or more. Further, the content of B is preferably 0.0025% or less.
Ｏ：０．００３０％以下
Ｏは、含有量が多くなると清浄度を低下させる。このため本発明ではできるだけ低減することが望ましい。特に、Ｏ含有量が０．００３０％を超えるとＣａＯ系介在物が粗大化して母材靭性を低下させてしまうため、０．００３０％以下とする。O: 0.0030% or less O decreases in cleanliness as the content increases. For this reason, it is desirable to reduce as much as possible in the present invention. In particular, when the O content exceeds 0.0030%, CaO-based inclusions become coarse and lower the base material toughness, so the content is made 0.0030% or less.
０＜［（Ｃａ−（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５］／Ｓ＜１
但し、Ｃａ、Ｏ、Ｓは、各元素の含有量（質量％）を表す。
本パラメータ式はＣａＳ上にＭｎＳが析出した複合硫化物の形態とするため、鋼中のＣａ、Ｓ、Ｏの含有量を規定するものである。
本パラメータ式の値が、０超え、１未満の場合、鋼を溶製する際の凝固段階でＣａＳが晶出し、ＣａＳの晶出後に固溶Ｓ量が確保されて、ＣａＳの表面上にＭｎＳが析出する。
ＭｎＳはフェライト核生成能を有し、その周囲にＭｎの希薄帯を形成してフェライト変態を促進し、溶接熱影響部の靭性を向上させる。ＭｎＳ上にＴｉＮ、ＢＮ、ＡｌＮ等のフェライト生成核が析出することによって、より一層、フェライト変態が促進される。
本パラメータ式の値が、０以下の場合には、ＣａＳが晶出せず、ＳはＭｎＳ単独の形態で析出し、溶接熱影響部において複合硫化物を微細分散させることができない。
一方、本パラメータ式の値が１以上の場合には、ＳがＣａによって完全に固定され、フェライト生成核として作用するＭｎＳが、ＣａＳ上に析出しないため、溶接熱影響部において複合硫化物を微細分散させることができない。
なお、本発明では、ＣａをＣａＳとして晶出させるために、Ｃａと結合力の強いＯ量をＣａ添加前に低減させておくことが必要であり、Ｃａ添加前の残存酸素量は、０．００３０％以下であることが好ましい。残存酸素量の低減方法としては、脱ガスを強化する、あるいは、脱酸剤を投入する、などの方法をとることができる。0 <[(Ca− (0.18 + 130 × Ca) × O) /1.25] / S <1
However, Ca, O, and S represent content (mass%) of each element.
This parameter formula defines the content of Ca, S, and O in the steel in order to form a composite sulfide in which MnS is precipitated on CaS.
When the value of this parameter formula is greater than 0 and less than 1, CaS crystallizes in the solidification stage when melting steel, and the amount of solid solution S is secured after crystallization of CaS, and MnS is formed on the surface of CaS. Precipitates.
MnS has a ferrite nucleation ability, and forms a Mn dilute band around it to promote ferrite transformation and improve the toughness of the weld heat affected zone. Ferrite transformation is further promoted by precipitation of ferrite-forming nuclei such as TiN, BN, and AlN on MnS.
When the value of this parameter formula is 0 or less, CaS does not crystallize, S precipitates in the form of MnS alone, and the composite sulfide cannot be finely dispersed in the weld heat affected zone.
On the other hand, when the value of this parameter formula is 1 or more, S is completely fixed by Ca, and MnS acting as a ferrite nucleation does not precipitate on CaS. Cannot be dispersed.
In the present invention, in order to crystallize Ca as CaS, it is necessary to reduce the amount of O having a strong binding force with Ca before the addition of Ca. It is preferable that it is 0030% or less. As a method for reducing the amount of residual oxygen, a method such as enhancing degassing or introducing a deoxidizer can be employed.
０．３６≦Ｃｅｑ≦０．５０
板厚５０ｍｍを超える厚鋼板の強度および集合組織強度を保つためには、Ｃｅｑ＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５で定義されるＣｅｑを０．３６〜０．５０とする。ただし、前記式における各元素記号は各元素の含有量（質量％）を表し、含有しない元素は０とする。好ましくはＣｅｑが０．３８〜０．４８である。0.36 ≦ Ceq ≦ 0.50
In order to maintain the strength and texture strength of the thick steel plate exceeding 50 mm, Ceq defined by Ceq = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 is set to 0.36 to 0.50. However, each element symbol in the above formula represents the content (% by mass) of each element, and the element not contained is 0. Preferably, Ceq is 0.38 to 0.48.
以上が本発明の基本成分組成で残部は、Ｆｅ及び不可避的不純物である。さらに、特性を向上させるため、Ｎｂ、Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏ、Ｖ、Ｍｇ、Ｚｒ、ＲＥＭの一種または二種以上を含有することが可能である。 The above is the basic component composition of the present invention, and the balance is Fe and inevitable impurities. Furthermore, in order to improve the characteristics, it is possible to contain one or more of Nb, Cu, Ni, Cr, Mo, V, Mg, Zr, and REM.
Ｎｂ：０．００３〜０．０４０％
Ｎｂは、ＮｂＣとしてフェライト変態時あるいは再加熱時に析出し、高強度化に寄与する。また、オーステナイト域の圧延において未再結晶域を拡大させる効果を有し、フェライトの細粒化に寄与するので、靭性の改善にも有効である。その効果は０．００３％以上の含有により発揮されるが、０．０４０％を超えて含有すると、粗大なＮｂＣが析出して逆に、靭性の低下を招くので、Ｎｂを含有する場合、その上限は０．０４０％とするのが好ましい。Ｎｂの含有量は、より好ましくは０．０１０％以上である。また、Ｎｂの含有量は、より好ましくは０．０３０％以下である。Nb: 0.003-0.040%
Nb precipitates as NbC during ferrite transformation or reheating, and contributes to increasing the strength. In addition, it has the effect of expanding the non-recrystallized region in the rolling of the austenite region and contributes to the refinement of ferrite, so it is also effective in improving toughness. The effect is exhibited by the content of 0.003% or more, but if it exceeds 0.040%, coarse NbC precipitates and conversely causes a decrease in toughness. The upper limit is preferably 0.040%. The Nb content is more preferably 0.010% or more. Further, the Nb content is more preferably 0.030% or less.
Ｃｕ：０．０１〜０．５％、Ｎｉ：０．０１〜２．０％
ＣｕおよびＮｉは、鋼の焼入れ性を高める元素である。圧延後の強度向上に直接寄与するとともに、靭性、高温強度、あるいは耐候性などの機能向上のために含有させることができ、これらの効果は、いずれも、０．０１％以上の含有によって発揮されるものの、過度の含有は靭性や溶接性を劣化させる。また、合金のコストも高くなってしまうため、Ｃｕ及び／又はＮｉを含有する場合には、それぞれの範囲を、Ｃｕは０．０１〜０．５％、Ｎｉは０．０１〜２．０％とする。Cu: 0.01 to 0.5%, Ni: 0.01 to 2.0%
Cu and Ni are elements that enhance the hardenability of steel. It contributes directly to strength improvement after rolling, and can be contained for improving functions such as toughness, high-temperature strength, or weather resistance. All of these effects are exhibited by inclusion of 0.01% or more. However, excessive inclusion deteriorates toughness and weldability. Moreover, since the cost of an alloy will also become high, when it contains Cu and / or Ni, Cu is 0.01 to 0.5%, Ni is 0.01 to 2.0%, respectively. And
Ｃｒ：０．０１〜０．５％、Ｍｏ：０．０１〜０．５％
ＣｒおよびＭｏは、いずれも鋼の焼入れ性を高める元素である。圧延後の強度向上に直接寄与するとともに、靭性、高温強度、あるいは耐候性などの機能向上のために含有させることができ、これらの効果は、いずれも、０．０１％以上の含有によって発揮されるものの、過度の含有は靭性や溶接性を劣化させる。靭性や溶接性を劣化させない範囲としては、Ｃｒ及び／又はＭｏを含有する場合のそれぞれの範囲を、０．０１〜０．５％とすることが好ましい。Cr: 0.01-0.5%, Mo: 0.01-0.5%
Cr and Mo are both elements that enhance the hardenability of steel. It contributes directly to strength improvement after rolling, and can be contained for improving functions such as toughness, high-temperature strength, or weather resistance. All of these effects are exhibited by inclusion of 0.01% or more. However, excessive inclusion deteriorates toughness and weldability. As a range which does not deteriorate toughness and weldability, it is preferable that each range in the case of containing Cr and / or Mo is 0.01 to 0.5%.
Ｖ：０．００１〜０．１０％
Ｖは、Ｖ（ＣＮ）として析出する析出強化によって、鋼の強度を向上させる元素であり、この効果はＶを０．００１％以上含有させることにより発揮される。しかし、Ｖを０．１０％を超えて含有すると、靭性を低下させる。このため、Ｖを含有させる場合には、Ｖの含有量を０．００１〜０．１０％の範囲とすることが好ましい。V: 0.001 to 0.10%
V is an element that improves the strength of the steel by precipitation strengthening that precipitates as V (CN), and this effect is exhibited by containing V in an amount of 0.001% or more. However, when V is contained exceeding 0.10%, toughness is reduced. For this reason, when it contains V, it is preferable to make content of V into 0.001 to 0.10% of range.
Ｍｇ：０．０００５〜０．００５０％、Ｚｒ：０．０００５〜０．０２００％、ＲＥＭ：０．０００５〜０．０２００％
Ｍｇ、ＺｒおよびＲＥＭ（希土類金属）はいずれも、酸化物の分散による靱性改善効果を有する元素である。このような効果を発現させるには、Ｍｇ、Ｚｒ、ＲＥＭのうちの少なくとも１種を０．０００５％以上含有させることが好ましい。一方、Ｍｇは０．００５０％超、ＺｒおよびＲＥＭは０．０２００％超を添加しても、その効果は飽和するだけである。よって、これらの元素を含有する場合は、Ｍｇ、Ｚｒ、ＲＥＭをそれぞれ、Ｍｇ：０．０００５〜０．００５０％、Ｚｒ：０．０００５〜０．０２００％、ＲＥＭ：０．０００５〜０．０２００％の範囲とすることが好ましい。Mg: 0.0005 to 0.0050%, Zr: 0.0005 to 0.0200%, REM: 0.0005 to 0.0200%
Mg, Zr, and REM (rare earth metals) are all elements that have an effect of improving toughness due to oxide dispersion. In order to express such an effect, it is preferable to contain 0.0005% or more of at least one of Mg, Zr, and REM. On the other hand, even if Mg exceeds 0.0050% and Zr and REM add more than 0.0200%, the effect is only saturated. Therefore, when these elements are contained, Mg, Zr, and REM are Mg: 0.0005 to 0.0050%, Zr: 0.0005 to 0.0200%, and REM: 0.0005 to 0.0200, respectively. % Is preferable.
上記以外の残部はＦｅ及び不可避的不純物とする。不可避的不純物としては、例えば、Ｓｂ、Ｓｎ、Ｚｎ、Ｃｏ等が挙げられ、これらの含有量の許容範囲としては、Ｓｂ：０．０１％以下、Ｓｎ：０．０１％以下、Ｚｎ：０．０１％以下、Ｃｏ：０．１％以下の範囲である。 The balance other than the above is Fe and inevitable impurities. Inevitable impurities include, for example, Sb, Sn, Zn, Co, etc. The allowable ranges of these contents are Sb: 0.01% or less, Sn: 0.01% or less, Zn: 0. The range is 01% or less and Co: 0.1% or less.
［鋼板内部の集合組織］
本発明では、圧延方向または圧延直角方向など板面に平行な方向に伝播する亀裂に対して亀裂伝播停止特性を向上させるために、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比および（２２２）面Ｘ線強度比、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比を規定する。なお、本発明における鋼板表面の板面（圧延面）に平行な面とは、鋼板の単純な表面だけではなく、鋼板表面をＸ線強度比（結晶面の集積度）が測定可能な面に処理した後の面を含む。例えば、鋼板の最表面がスケールで覆われている時などは、それを取り除いた面を言う。また、鋼板の最表面が鏡面になっており、そのままＸ線強度比（結晶面の集積度）を測定できる場合などは、鋼板の表面（圧延面）そのものをいう。[A texture inside the steel sheet]
In the present invention, the (211) plane X in the plane parallel to the plate surface at the center of the plate thickness is used in order to improve the crack propagation stop characteristic for cracks propagating in the direction parallel to the plate surface such as the rolling direction or the direction perpendicular to the rolling direction. A line intensity ratio, a (222) plane X-ray intensity ratio, and a (211) plane X-ray intensity ratio in a plane parallel to the plate surface of the steel sheet surface are defined. In addition, the surface parallel to the plate surface (rolled surface) of the steel plate surface in the present invention is not only a simple surface of the steel plate but also a surface on which the X-ray intensity ratio (degree of integration of crystal planes) can be measured. Including the surface after processing. For example, when the outermost surface of a steel plate is covered with a scale, it refers to the surface from which it has been removed. In addition, when the outermost surface of the steel plate is a mirror surface and the X-ray intensity ratio (degree of crystal plane integration) can be measured as it is, the surface of the steel plate (rolled surface) itself is referred to.
板厚中央および鋼板表面において（２１１）面を発達させると同時に、板厚中央における（２２２）面の発達を抑制することで、脆性亀裂が直進的に進展することが阻害され、その結果、脆性亀裂伝播停止特性の向上が可能となる。 By developing the (211) plane at the center of the plate thickness and the surface of the steel sheet, and simultaneously suppressing the development of the (222) plane at the center of the plate thickness, the brittle cracks are prevented from progressing in a straight line. The crack propagation stop characteristic can be improved.
具体的に、最近のコンテナ船やバルクキャリアーなど船体外板に用いられるようになった板厚５０ｍｍ以上の厚肉材で、構造安全性を確保する上で目標とされるＫｃａ（−１０℃）≧６０００Ｎ／ｍｍ３／２の脆性亀裂伝播停止特性を得る場合には、鋼板の板厚中央（板厚１／２位置を意味する。なお、Ｘ線強度比（結晶面の集積度）の測定においては、板厚１／２の位置から上下方向に板厚の数％の位置誤差［板厚の０％超〜５％以下の位置誤差］は許容される。）の板面に平行な面における（２１１）面Ｘ線強度比を１．５以上かつ（２２２）面Ｘ線強度比を２．５以下、および鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比を１．２以上とする必要がある。より優れた亀裂伝播停止性能が要求される場合は、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比を１．６以上、（２２２）面Ｘ線強度比を２．２以下、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比を１．４以上とすることが好ましい。Specifically, Kca (-10 ° C), which is the target for ensuring structural safety, is a thick material with a thickness of 50mm or more that has come to be used for hull outer plates such as recent container ships and bulk carriers. In order to obtain a brittle crack propagation stop characteristic of ≧ 6000 N / mm 3/2, the center of the plate thickness (meaning the plate thickness 1/2 position. Note that the X-ray intensity ratio (the degree of crystal plane integration) is measured. In this case, a position error of several percent of the plate thickness in the vertical direction from the position of the plate thickness 1/2 (position error of more than 0% to 5% or less of the plate thickness) is allowed. The (211) plane X-ray intensity ratio is 1.5 or more, the (222) plane X-ray intensity ratio is 2.5 or less, and the (211) plane X-ray intensity ratio in a plane parallel to the plate surface of the steel sheet surface is 1 Need to be 2 or more. When better crack propagation stopping performance is required, the (211) plane X-ray intensity ratio is 1.6 or more and the (222) plane X-ray intensity ratio is 2. It is preferable that the (211) plane X-ray intensity ratio in a plane parallel to the plate surface of the steel plate surface is 2 or more and 1.4 or more.
なお、ここで、（２１１）面Ｘ線強度比とは、対象材の（２１１）結晶面の集積度を表す数値で、対象材の（２１１）反射のＸ線回折強度（Ｉ（２１１））と、集合組織のないランダムな標準試料の（２１１）反射のＸ線回折強度（Ｉ０（２１１））との比（Ｉ（２１１）／Ｉ０（２１１））を指す。（２２２）面Ｘ線強度比についても同様である。すなわち、対象材の（２２２）反射のＸ線回折強度（Ｉ（２２２））と、集合組織のないランダムな標準試料の（２２２）反射のＸ線回折強度（Ｉ０（２２２））との比（Ｉ（２２２）／Ｉ０（２２２））を指す。Here, the (211) plane X-ray intensity ratio is a numerical value indicating the degree of integration of the (211) crystal plane of the target material, and the (211) reflection X-ray diffraction intensity (I (211)) of the target material. And the ratio (I (211) / I 0 (211)) of X-ray diffraction intensity (I 0 (211)) of (211) reflection of a random standard sample without texture. The same applies to the (222) plane X-ray intensity ratio. That is, the ratio between the (222) reflection X-ray diffraction intensity (I (222)) of the target material and the (222) reflection X-ray diffraction intensity (I 0 (222)) of the random standard sample without texture. (I (222) / I 0 (222)).
以下、板厚５０ｍｍを超える厚肉材で、Ｋｃａ（−１０℃）≧６０００Ｎ／ｍｍ３／２の脆性亀裂伝播停止特性を得る場合に好ましい製造条件について説明する。Hereinafter, preferable manufacturing conditions will be described when a brittle crack propagation stop characteristic of Kca (−10 ° C.) ≧ 6000 N / mm 3/2 is obtained with a thick material exceeding a plate thickness of 50 mm.
［製造条件］
上記成分組成になる溶鋼を、転炉等で溶製し、連続鋳造等で鋼素材（スラブ）とし、１０００〜１２００℃に加熱後、熱間圧延を行う。[Production conditions]
The molten steel having the above component composition is melted in a converter or the like, made into a steel material (slab) by continuous casting or the like, heated to 1000 to 1200 ° C., and then hot-rolled.
加熱温度が１０００℃未満では、オーステナイト再結晶温度域における圧延を行う時間が十分に確保できない一方で、１２００℃超では、オーステナイト粒が粗大化し、靭性の低下を招くばかりか、酸化ロスが顕著となって、歩留が低下する。従って、鋼素材の加熱温度は、１０００〜１２００℃の範囲とする。鋼板の靭性向上の観点から好ましい加熱温度の範囲は１０００〜１１５０℃であり、さらに好ましくは１０３０〜１１３０℃である。 When the heating temperature is less than 1000 ° C., sufficient time for rolling in the austenite recrystallization temperature range cannot be secured, but when it exceeds 1200 ° C., the austenite grains become coarse, leading to a decrease in toughness, and oxidation loss is remarkable. As a result, the yield decreases. Therefore, the heating temperature of the steel material is in the range of 1000 to 1200 ° C. The range of preferable heating temperature from a viewpoint of the toughness improvement of a steel plate is 1000-1150 degreeC, More preferably, it is 1030-1130 degreeC.
次いで、熱間圧延を施す。
熱間圧延はまず、板厚中央の温度がオーステナイト再結晶温度域で圧延を行い、ついで板厚中央の温度がオーステナイト未再結晶温度域で圧延を行う。このとき、圧延途中で鋼板の表裏面から加熱を行い板厚方向の温度分布を制御することにより、鋼板内部と鋼板の表裏面の温度差を小さくする制御を行う。これにより目標とする板厚中央と鋼板表面における集合組織強度が得られる。この制御は、前記オーステナイト再結晶温度域での圧延終了後に行われるのが好ましい。より好ましくは、前記オーステナイト再結晶温度域での圧延終了後、オーステナイト未再結晶温度域での圧延開始前である。Next, hot rolling is performed.
In the hot rolling, first, rolling is performed at a temperature at the center of the sheet thickness in the austenite recrystallization temperature range, and then, rolling is performed at a temperature in the center of the sheet thickness at the austenite non-recrystallization temperature range. At this time, control is performed to reduce the temperature difference between the inside of the steel sheet and the front and back surfaces of the steel sheet by controlling the temperature distribution in the thickness direction by heating from the front and back surfaces of the steel sheet during rolling. As a result, target texture strength at the center of the plate thickness and the surface of the steel plate can be obtained. This control is preferably performed after the end of rolling in the austenite recrystallization temperature range. More preferably, after completion of rolling in the austenite recrystallization temperature range, before rolling in the austenite non-recrystallization temperature range.
また、その際、オーステナイト未再結晶温度域での圧延が、鋼板表面の温度−板厚中央の温度≧−４０℃となる条件で行われるように前記制御が行われることが好ましい。これにより、板厚中央をより低い温度で圧延しつつ、鋼板表面にフェライトが生成することを抑制することができる。より好ましくは鋼板表面の温度−板厚中央の温度≧−２０℃での圧延である。また、前記条件でのオーステナイト未再結晶温度域での圧延は、当該オーステナイト未再結晶温度域での圧延を開始してから終了するまでの間の少なくとも一部で行われることが好ましい。例えば、オーステナイト未再結晶温度域での圧延を開始してから終了するまでの間、前記条件を保持して圧延を行ってもよいし、オーステナイト未再結晶温度域での圧延の開始時、終了時、或いはこれらの間の圧延の一部（一時点）において前記条件による圧延を行なってもよい。好ましくは、オーステナイト未再結晶温度域での圧延終了時において前記条件が満たされているように圧延を行う。オーステナイト未再結晶温度域での圧延終了時に前記条件が満たされていることで鋼板の特性がより高められる。
また、好ましくは、上記制御をオーステナイト未再結晶温度域での圧延開始前に行い、オーステナイト未再結晶温度域での圧延開始時から、或いは、前記圧延開始時より一定時間経過後から、前記条件による圧延を行う。より好ましくは、上記制御をオーステナイト未再結晶温度域での圧延開始前に行い、オーステナイト未再結晶温度域での圧延開始時から、或いは、前記圧延開始時より一定時間経過後から、前記圧延終了時までの間、前記条件による圧延を行う。
なお、前記鋼板の表裏面から加熱を行うための手段としては、特に制限されないが、例えば、雰囲気炉による加熱や、高周波による加熱等が挙げられる。また、前記鋼板表面の温度−板厚中央の温度の上限は、特に制限されないが、母材靭性の確保および材質均一性の観点からは、８０℃以下が好ましく、６０℃以下がより好ましい。At that time, it is preferable that the control is performed so that rolling in the austenite non-recrystallization temperature region is performed under the condition that the temperature of the steel sheet surface-the temperature at the center of the plate thickness ≧ −40 ° C. Thereby, it can suppress that a ferrite produces | generates on the steel plate surface, rolling a plate | board thickness center at lower temperature. More preferably, the rolling is performed at a temperature of the steel sheet surface-a temperature at the center of the sheet thickness? Moreover, it is preferable that the rolling in the austenite non-recrystallization temperature region is performed at least partly from the start to the end of rolling in the austenite non-recrystallization temperature region. For example, during the period from the start to the end of rolling in the austenite non-recrystallization temperature range, the above-mentioned conditions may be maintained and the rolling may be performed. Rolling according to the above conditions may be performed at the time or part of the rolling (temporary point) between them. Preferably, the rolling is performed so that the above condition is satisfied at the end of rolling in the austenite non-recrystallization temperature range. The properties of the steel sheet can be further improved by satisfying the above conditions at the end of rolling in the austenite non-recrystallization temperature region.
Preferably, the above control is performed before the start of rolling in the austenite non-recrystallization temperature range, from the start of rolling in the austenite non-recrystallization temperature range, or after a lapse of a certain time from the start of rolling. Roll by. More preferably, the above control is performed before the start of rolling in the austenite non-recrystallization temperature range, and from the start of rolling in the austenite non-recrystallization temperature range, or after a lapse of a certain time from the start of rolling, the rolling is completed. Until that time, rolling is performed under the above conditions.
The means for heating from the front and back surfaces of the steel sheet is not particularly limited, and examples thereof include heating by an atmospheric furnace and heating by high frequency. Moreover, the upper limit of the temperature of the steel sheet surface-the center of the plate thickness is not particularly limited, but is preferably 80 ° C. or less, and more preferably 60 ° C. or less, from the viewpoint of securing the base material toughness and material uniformity.
なお、オーステナイト未再結晶温度域での圧延は、板厚中央の温度がＡｒ３〜Ａｒ３＋３０℃の範囲での累積圧下率が４０％以上とすることが好ましい。より好ましくはＡｒ３〜Ａｒ３＋２０℃の範囲での累積圧下率が４０％以上であり、さらに好ましくはＡｒ３〜Ａｒ３＋２０℃の範囲での累積圧下率が５０％以上である。オーステナイト未再結晶温度域での圧延を終了する前に鋼板の表裏面から加熱し未結晶温度域での鋼板の板厚方向の温度分布を制御することで板厚中央の温度がＡｒ３点直上で圧延しても表層フェライトの生成を抑制することができる。ここで、Ａｒ３点（Ａｒ３変態点）は下記式を用いて算出した値を用いるものとする。
Ａｒ３点（℃）＝９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−１５Ｃｒ−８０Ｍｏ
ただし、上記式中のＣ、Ｓｉ、Ｍｎ、Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏは、前記各元素の含有量（質量％）を表し、含有しない元素は０とする。
また、上記オーステナイト再結晶温度域での圧延は、特に制限されないが、累積圧下率が１０％以上であることが好ましい。より好ましくは、累積圧下率が１５％以上である。In the rolling in the austenite non-recrystallization temperature range, it is preferable that the cumulative reduction ratio is 40% or more when the temperature at the center of the sheet thickness is in the range of Ar 3 to Ar 3 + 30 ° C. More preferably, the cumulative rolling reduction in the range of Ar 3 to Ar 3 + 20 ° C. is 40% or more, and still more preferably, the cumulative rolling reduction in the range of Ar 3 to Ar 3 + 20 ° C. is 50% or more. Before finishing rolling in the austenite non-recrystallization temperature range, heating from the front and back surfaces of the steel plate and controlling the temperature distribution in the thickness direction of the steel plate in the non-crystal temperature range, the temperature at the center of the plate thickness is just above the Ar 3 point. The production of surface ferrite can be suppressed even when rolled at. Here, the Ar 3 point (Ar 3 transformation point) is a value calculated using the following equation.
Ar 3 points (° C.) = 910-310C-80Mn-20Cu-55Ni-15Cr-80Mo
However, C, Si, Mn, Cu, Ni, Cr, and Mo in the above formulas represent the content (% by mass) of each of the above elements, and the elements that do not contain 0.
The rolling in the austenite recrystallization temperature range is not particularly limited, but the cumulative rolling reduction is preferably 10% or more. More preferably, the cumulative rolling reduction is 15% or more.
さらに、熱間圧延終了時における鋼板表面の温度と板厚中央の温度との温度差が５℃以内であることが好ましい。また、熱間圧延終了時における板厚中央の温度がＡｒ３℃〜（Ａｒ３+３０）℃であることが好ましい。これにより鋼板の特性がより高められる。Furthermore, it is preferable that the temperature difference between the temperature of the steel sheet surface at the end of hot rolling and the temperature at the center of the sheet thickness is within 5 ° C. Further, it is preferred that the temperature of the plate thickness center at the end of hot rolling is Ar 3 ℃ ~ (Ar 3 +30 ) ℃. Thereby, the characteristic of a steel plate is improved more.
圧延が終了した鋼板は、Ａｒ３点以上の温度から２℃／ｓ以上の冷却速度で６００℃以下まで冷却することが好ましい。冷却速度を２℃／ｓ以上とすることで、圧延時に発達させた集合組織強度を保つことができる。It is preferable that the rolled steel sheet is cooled from a temperature of 3 or higher Ar to 600 ° C. or lower at a cooling rate of 2 ° C./s or higher. By setting the cooling rate to 2 ° C./s or more, the texture strength developed during rolling can be maintained.
圧延および冷却後に焼戻処理を行う場合は、Ａｃ１点以下で行うことが好ましい。焼戻処理がＡｃ１点以上の場合には、圧延時に発達させた集合組織を失うおそれがあるからである。When performing a tempering process after rolling and cooling, it is preferable to carry out at Ac 1 point or less. This is because when the tempering treatment is at least one point of Ac, the texture developed during rolling may be lost.
なお、以上の説明において、板厚中央の温度は、放射温度計で測定した鋼板の表面温度からの伝熱計算より求める。また、圧延後の冷却条件における温度条件は、板厚中央の温度とする。本発明における高強度厚鋼板とは、引張り強さが５８０ＭＰａ以上の厚鋼板をいう。本発明における高強度厚鋼板は６００ＭＰａ以上の引張り強さを有することが好ましい。本発明における高強度厚鋼板の板厚は、５０ｍｍ超が好ましく、５５ｍｍ以上がより好ましく、６０ｍｍ以上がさらに好ましい。 In the above description, the temperature at the center of the plate thickness is obtained from the heat transfer calculation from the surface temperature of the steel plate measured with a radiation thermometer. The temperature condition in the cooling condition after rolling is the temperature at the center of the plate thickness. The high strength thick steel plate in the present invention refers to a thick steel plate having a tensile strength of 580 MPa or more. The high-strength thick steel plate in the present invention preferably has a tensile strength of 600 MPa or more. The thickness of the high-strength thick steel plate in the present invention is preferably more than 50 mm, more preferably 55 mm or more, and further preferably 60 mm or more.
また、本発明における高強度厚鋼板は、シャルピー衝撃試験による板厚中央の破面遷移温度が−６０℃以下（ｖＴｒｓ≦−６０℃）であることが好ましい。さらに、本発明における高強度厚鋼板は、シャルピー衝撃試験による鋼板表面から板厚方向に５ｍｍの位置の破面遷移温度が−６０℃以下（ｖＴｒｓ≦−６０℃）であることが好ましい。これにより、鋼板の特性がより高められる。なお、前記板厚中央の破面遷移温度および鋼板表面から板厚方向に５ｍｍの位置の破面遷移温度の要件を満たす高強度厚鋼板を得るには、上述した板厚中央の温度がＡｒ３〜Ａｒ３＋３０℃の範囲での累積圧下率を４５％以上とするか、または、板厚中央の温度がＡｒ３〜Ａｒ３＋３０℃の範囲での累積圧下率が４５％未満の場合には、熱間圧延終了時における板厚中央の温度を７４０℃以下とすることが好ましい。In the high-strength thick steel plate according to the present invention, the fracture surface transition temperature at the center of the plate thickness by the Charpy impact test is preferably −60 ° C. or lower (vTrs ≦ −60 ° C.). Furthermore, the high-strength thick steel plate in the present invention preferably has a fracture surface transition temperature of −60 ° C. or less (vTrs ≦ −60 ° C.) at a position 5 mm from the steel plate surface in the plate thickness direction by the Charpy impact test. Thereby, the characteristic of a steel plate is improved more. In order to obtain a high-strength thick steel plate that satisfies the requirements of the fracture surface transition temperature at the center of the plate thickness and the fracture surface transition temperature at a position 5 mm from the steel plate surface in the plate thickness direction, the temperature at the center of the plate thickness described above is Ar 3. When the cumulative rolling reduction in the range of ˜Ar 3 + 30 ° C. is 45% or more, or the cumulative rolling reduction in the range of the thickness center of Ar 3 to Ar 3 + 30 ° C. is less than 45% The temperature at the center of the plate thickness at the end of hot rolling is preferably 740 ° C. or lower.
次に、本発明の実施例について説明する。
表１に示す各成分組成の溶鋼を、転炉で溶製し、連続鋳造法で鋼素材とし、表２に示す加熱温度で加熱し、板厚５５〜１００ｍｍに熱間圧延後、冷却を行い厚鋼板を得た。表２に熱間圧延条件と冷却条件を示す。熱間圧延中に、鋼板の表裏面から加熱を行い板厚方向の温度分布制御を行ったものについては、表２の圧延中の加熱有無に○を示した。前記加熱は、オーステナイト再結晶温度域での圧延終了後、オーステナイト未再結晶温度域での圧延開始前に行い、前記加熱後、３０秒以内にオーステナイト未再結晶温度域での圧延を開始した。なお、前記加熱を行わなかったものについては、表２の圧延中の加熱有無に×を示した。熱間圧延後は直ちに板厚中央の冷却速度が２〜１０℃／ｓで３５０〜５００℃の範囲まで冷却し、その後、放冷した。なお、前記鋼板の表裏面からの加熱は、雰囲気炉加熱装置により行った。また、表１のＡｒ３は、前述したものと同様の計算式により求めた。Next, examples of the present invention will be described.
Molten steel of each component composition shown in Table 1 is melted in a converter, made into a steel material by a continuous casting method, heated at the heating temperature shown in Table 2, and cooled after hot rolling to a plate thickness of 55 to 100 mm. A thick steel plate was obtained. Table 2 shows hot rolling conditions and cooling conditions. In the case of heating from the front and back surfaces of the steel sheet during hot rolling and controlling the temperature distribution in the plate thickness direction, the presence or absence of heating during rolling in Table 2 was indicated as ◯. The heating was performed after the end of rolling in the austenite recrystallization temperature range and before the start of rolling in the austenite non-recrystallization temperature range. After the heating, rolling in the austenite non-recrystallization temperature range was started within 30 seconds. In addition, about what did not perform the said heating, x was shown in the presence or absence of the heating during rolling of Table 2. Immediately after the hot rolling, the sheet was cooled to a range of 350 to 500 ° C. at a cooling rate of 2 to 10 ° C./s at the center of the plate thickness, and then allowed to cool. In addition, the heating from the front and back surfaces of the steel plate was performed by an atmospheric furnace heating device. Further, Ar 3 in Table 1 was obtained by the same calculation formula as described above.
得られた厚鋼板について、板厚１／４の位置より、Φ１４のＪＩＳ １４Ａ号試験片を採取し、引張試験を行い、降伏強度（ＹＳ）、引張強さ（ＴＳ）を測定した。 About the obtained thick steel plate, the JIS14A test piece of (phi) 14 was extract | collected from the position of the plate | board thickness 1/4, the tensile test was done, and the yield strength (YS) and the tensile strength (TS) were measured.
また、板厚中央の破面遷移温度（ｖＴｒｓ）については、板厚１／２位置より、ＪＩＳ ４号衝撃試験片を試験片の長手軸の方向が圧延方向と平行となるように採取し、シャルピー衝撃試験を行って、−２０℃〜−１００℃におけるシャルピー吸収エネルギーおよびその脆性破面率を測定し、脆性−延性破面遷移温度(ｖＴｒｓ)（℃）を求めた。
鋼板表面から板厚方向に５ｍｍの位置（表３中に「表層部」と記載）の破面遷移温度（ｖＴｒｓ）については、鋼板の表面に形成されているスケール（黒皮）を除去した後、該鋼板の表面から、ＪＩＳ ４号衝撃試験片を試験片の長手軸の方向が圧延方向と平行となるように採取した。すなわち、前記試験片の厚さは１０ｍｍであるから、前記試験片における測定位置は、該試験片の厚さ方向の中心位置、すなわち、鋼板表面から板厚方向に５ｍｍの位置となる。この試験片に対してシャルピー衝撃試験を行って、−２０℃〜−１００℃におけるシャルピー吸収エネルギーおよびその脆性破面率を測定し、鋼板表面から板厚方向に５ｍｍの位置の脆性−延性破面遷移温度(ｖＴｒｓ)（℃）を求めた。Also, for the fracture surface transition temperature (vTrs) at the center of the plate thickness, from the position of the plate thickness 1/2, a JIS No. 4 impact test piece was collected so that the direction of the longitudinal axis of the test piece was parallel to the rolling direction, A Charpy impact test was performed to measure Charpy absorption energy at −20 ° C. to −100 ° C. and its brittle fracture surface ratio, and a brittle-ductile fracture surface transition temperature (vTrs) (° C.) was determined.
For the fracture surface transition temperature (vTrs) at a position 5 mm from the steel sheet surface in the thickness direction (described as “surface layer” in Table 3), after removing the scale (black skin) formed on the steel sheet surface From the surface of the steel plate, a JIS No. 4 impact test piece was collected so that the direction of the longitudinal axis of the test piece was parallel to the rolling direction. That is, since the thickness of the test piece is 10 mm, the measurement position on the test piece is the center position in the thickness direction of the test piece, that is, the position of 5 mm from the steel plate surface to the plate thickness direction. A Charpy impact test was performed on this test piece to measure Charpy absorbed energy at −20 ° C. to −100 ° C. and its brittle fracture surface ratio, and a brittle-ductile fracture surface at a position of 5 mm in the thickness direction from the steel sheet surface. The transition temperature (vTrs) (° C.) was determined.
また、厚鋼板の集合組織を評価するため、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比と（２２２）面Ｘ線強度比、および鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比を測定した。
測定方法はまず、板厚表層下０．５ｍｍあるいは板厚中央から板厚１ｍｍのサンプルを採取し、板面に平行な面を機械研磨・電解研磨することにより、Ｘ線回折用の試験片を用意する。なお、板厚表層部の場合には、最表面に近い方の面を研磨するものとする。この試験片を用いて、Ｍｏ線源を用いてＸ線回折装置を使用して、Ｘ線回折測定を実施し、（２１１）面Ｘ線強度比および（２２２）面Ｘ線強度比を求めた。Further, in order to evaluate the texture of the thick steel plate, the (211) plane X-ray intensity ratio and the (222) plane X-ray intensity ratio in a plane parallel to the plate surface at the center of the plate thickness, and the plate surface of the steel plate surface are parallel. The (211) plane X-ray intensity ratio in the plane was measured.
The measurement method is as follows. First, a sample with a surface thickness of 0.5 mm or a thickness of 1 mm is taken from the center of the plate thickness, and a surface parallel to the plate surface is mechanically polished / electropolished to prepare a test piece for X-ray diffraction. prepare. In the case of the plate thickness surface layer portion, the surface closest to the outermost surface is polished. Using this test piece, X-ray diffraction measurement was performed using an X-ray diffractometer using a Mo ray source, and a (211) plane X-ray intensity ratio and a (222) plane X-ray intensity ratio were obtained. .
次に、脆性亀裂伝播停止特性を評価するため、温度勾配型ＥＳＳＯ試験を行い、−１０℃におけるＫｃａ値（以下、Ｋｃａ（−１０℃）Ｎ／ｍｍ３／２とも記す）を求めた。Ｋｃａ（−１０℃）Ｎ／ｍｍ３／２が６０００Ｎ／ｍｍ３／２以上であれば脆性亀裂伝播停止特性に優れると評価できる。Next, in order to evaluate the brittle crack propagation stop characteristic, a temperature gradient type ESSO test was performed, and a Kca value at −10 ° C. (hereinafter also referred to as Kca (−10 ° C.) N / mm 3/2 ) was obtained. Kca (-10 ℃) N / mm 3/2 can be evaluated with excellent brittle crack propagation stop characteristics if 6000 N / mm 3/2 or more.
さらに、各厚鋼板から採取した継手用試験板に、Ｖ開先を施し、エレクトロガスアーク溶接により大入熱溶接継手を作製した。この際の入熱量（ｋＪ／ｃｍ）を表３に示す。また、得られた大入熱溶接継手から切欠位置をボンド部とするＪＩＳ４号衝撃試験片を採取し、試験温度−４０℃でシャルピー衝撃試験を実施し、同一条件で実施した試験片３本の吸収エネルギーの平均値を吸収エネルギーｖＥ−４０（Ｊ）として求めた。この吸収エネルギーｖＥ−４０（Ｊ）が８０Ｊ以上であれば前記継手の靭性に優れると評価できる。Furthermore, V groove was given to the test plate for joint extract | collected from each thick steel plate, and the high heat input welded joint was produced by electrogas arc welding. The heat input (kJ / cm) at this time is shown in Table 3. In addition, a JIS No. 4 impact test piece having a notch position as a bond portion was taken from the obtained high heat input welded joint, and a Charpy impact test was performed at a test temperature of −40 ° C. The average value of absorbed energy was determined as absorbed energy vE- 40 (J). If this absorbed energy vE- 40 (J) is 80J or more, it can be evaluated that the joint has excellent toughness.
表３にこれらの試験結果を示す。 Table 3 shows the results of these tests.
表３に示された結果から、本発明例では、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比が１．５以上、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比が１．２以上であり、板厚中央の板面に平行な面における（２２２）面Ｘ線強度比が２．５以下であるとともに、Ｋｃａ（−１０℃）Ｎ／ｍｍ３／２が６０００Ｎ／ｍｍ３／２以上と優れた脆性亀裂伝播停止特性が得られ、さらに溶接継手ボンド部のシャルピー吸収エネルギーが８０Ｊ以上と優れた継手靱性を有することが分かる。From the results shown in Table 3, in the example of the present invention, the (211) plane X-ray intensity ratio in the plane parallel to the plate surface at the center of the plate thickness is 1.5 or more, and in the plane parallel to the plate surface of the steel plate surface ( 211) The plane X-ray intensity ratio is 1.2 or more, the (222) plane X-ray intensity ratio in a plane parallel to the plate surface at the center of the plate thickness is 2.5 or less, and Kca (−10 ° C.) N It can be seen that an excellent brittle crack propagation stop characteristic is obtained when / mm 3/2 is 6000 N / mm 3/2 or more, and that the Charpy absorbed energy of the weld joint bond portion is 80 J or more and has excellent joint toughness.
一方、本発明を外れる厚鋼板Ｎｏ.３、５、６、２１〜３１の場合、成分組成、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比、（２２２）面Ｘ線強度比、鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比、引張り強さのいずれか１つ以上が本発明の規定を満たしておらず、Ｋｃａ（−１０℃）Ｎ／ｍｍ３／２、吸収エネルギーｖＥ−４０（Ｊ）のいずれか１つ以上で必要な特性が得られていない。具体的には、Ｋｃａ（−１０℃）が６０００Ｎ／ｍｍ３／２未満となるか、ｖＥ−４０（Ｊ）が８０Ｊ未満となり、必要な特性が得られていない。On the other hand, in the case of thick steel plates No. 3, 5, 6, 21 to 31 that depart from the present invention, the component composition, the (211) plane X-ray intensity ratio in the plane parallel to the plane at the center of the thickness, the (222) plane X Any one or more of the line strength ratio, the (211) plane X-ray strength ratio and the tensile strength in a plane parallel to the plate surface of the steel sheet does not satisfy the provisions of the present invention, and Kca (−10 ° C.) N Necessary characteristics are not obtained with any one or more of / mm 3/2 and absorbed energy vE- 40 (J). Specifically, Kca (−10 ° C.) is less than 6000 N / mm 3/2 or vE −40 (J) is less than 80 J, and necessary characteristics are not obtained.
Claims (6)
Ｃ：０．０３〜０．２０％、
Ｓｉ：０．０１〜０．３０％、
Ｍｎ：１．５〜３．０％、
Ｐ：０．０２％以下、
Ｓ：０．０００５〜０．０１％、
Ｔｉ：０．００５〜０．０３０％、
Ａｌ：０．００５〜０．０８０％、
Ｎ：０．００２５〜０．００７５％、
Ｃａ：０．０００３〜０．００３０％、
Ｂ：０．０００３〜０．００３０％、
Ｏ：０．００３０％以下を含有し、かつ、Ｃａ、Ｏ、Ｓが下記（１）式を満足し、下記（２）式で定義されるＣｅｑが０．３６〜０．５０の範囲にあり、残部がＦｅおよび不可避的不純物からなり、
鋼板表面の板面に平行な面における（２１１）面Ｘ線強度比が１．２以上、板厚中央の板面に平行な面における（２１１）面Ｘ線強度比が１．５以上かつ（２２２）面Ｘ線強度比が２．５以下であり、
３００ｋＪ／ｃｍ超の入熱量で溶接した溶接継手ボンド部の−４０℃におけるシャルピー吸収エネルギー（ｖＥ−４０）が８０Ｊ以上であり、ＥＳＳＯ試験による−１０℃におけるＫｃａ値（Ｋｃａ（−１０℃））が６０００Ｎ／ｍｍ３／２以上であり、引張り強さが５８０ＭＰａ以上であり、板厚が５０ｍｍ超である、高強度厚鋼板。
０＜［（Ｃａ−（０．１８＋１３０×Ｃａ）×Ｏ）／１．２５］／Ｓ＜１ ・・・（１）
Ｃｅｑ＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５ ・・・（２）
ただし、上記（１）式および（２）式における各元素記号は各元素の含有量（質量％）を表し、含有しない元素は０とする。Ingredient composition is mass%,
C: 0.03 to 0.20%
Si: 0.01-0.30%,
Mn: 1.5-3.0%
P: 0.02% or less,
S: 0.0005 to 0.01%
Ti: 0.005 to 0.030%,
Al: 0.005 to 0.080%,
N: 0.0025 to 0.0075%,
Ca: 0.0003 to 0.0030%,
B: 0.0003 to 0.0030%,
O: 0.0030% or less, and Ca, O, and S satisfy the following formula (1), and Ceq defined by the following formula (2) is in the range of 0.36 to 0.50. The balance consists of Fe and inevitable impurities,
The (211) plane X-ray intensity ratio in the plane parallel to the plate surface of the steel plate surface is 1.2 or more, and the (211) plane X-ray intensity ratio in the plane parallel to the plate surface at the center of the plate thickness is 1.5 or more ( 222) the plane X-ray intensity ratio is 2.5 or less,
The Charpy absorbed energy (vE -40 ) at -40 ° C of the welded joint part welded with a heat input of more than 300 kJ / cm is 80 J or more, and the Kca value (Kca (-10 ° C)) at -10 ° C by the ESSO test. There is a 6000 N / mm 3/2 or more and tensile strength of 580MPa or more, the thickness is 50mm greater, high strength thick steel plate.
0 <[(Ca− (0.18 + 130 × Ca) × O) /1.25] / S <1 (1)
Ceq = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (2)
However, each element symbol in the above formulas (1) and (2) represents the content (mass%) of each element, and the elements not contained are 0.
Ｎｂ：０．００３〜０．０４０％、
Ｃｕ：０．０１〜０．５％、
Ｎｉ：０．０１〜２．０％、
Ｃｒ：０．０１〜０．５％、
Ｍｏ：０．０１〜０．５％のなかから選ばれる１種または２種以上を含有する、請求項１〜３のいずれか一項に記載の高強度厚鋼板。Ingredient composition is further mass%,
Nb: 0.003-0.040%,
Cu: 0.01 to 0.5%,
Ni: 0.01 to 2.0%,
Cr: 0.01 to 0.5%
Mo: The high-strength thick steel plate as described in any one of Claims 1-3 containing 1 type, or 2 or more types chosen from 0.01-0.5%.
Ｖ：０．００１〜０．１０％、
Ｍｇ：０．０００５〜０．００５０％、
Ｚｒ：０．０００５〜０．０２００％、
ＲＥＭ：０．０００５〜０．０２００％のなかから選ばれる１種または２種以上を含有する、請求項１〜４のいずれか一項に記載の高強度厚鋼板。Ingredient composition is further mass%,
V: 0.001 to 0.10%,
Mg: 0.0005 to 0.0050%,
Zr: 0.0005 to 0.0200%,
The high-strength thick steel plate according to any one of claims 1 to 4, comprising one or more selected from REM: 0.0005 to 0.0200%.
熱間圧延終了時における鋼板表面の温度と板厚中央の温度との温度差を５℃以内とし、かつ、板厚中央の温度をＡｒ３℃〜（Ａｒ３+３０）℃とする、高強度厚鋼板の製造方法。It is a manufacturing method of the high intensity | strength thick steel plate as described in any one of Claims 1-5, Comprising: After heating the steel raw material which has the said component composition to the temperature of 1000-1200 degreeC, a sheet thickness center is austenite. Hot rolling is performed in the recrystallization temperature range, and then in the austenite non-recrystallization temperature range, and heating is performed from the front and back surfaces of the steel sheet during the hot rolling, and at least hot rolling in the austenite non-recrystallization temperature range. The temperature distribution in the plate thickness direction is controlled so as to perform a part of the temperature under the condition that the temperature of the steel plate surface−the temperature at the center of the plate thickness ≧ −40 ° C.,
High strength at which the temperature difference between the surface temperature of the steel sheet at the end of hot rolling and the temperature at the center of the sheet thickness is within 5 ° C., and the temperature at the center of the sheet thickness is Ar 3 to (Ar 3 +30) ° C. Manufacturing method of thick steel plate.
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