JPWO2009072559A1 - Manufacturing method of thick high-strength steel plate with excellent brittle fracture propagation stop characteristics and high heat input weld heat affected zone toughness, and thick high strength steel plate with excellent brittle fracture propagation stop characteristics and high heat input weld heat affected zone toughness - Google Patents

Manufacturing method of thick high-strength steel plate with excellent brittle fracture propagation stop characteristics and high heat input weld heat affected zone toughness, and thick high strength steel plate with excellent brittle fracture propagation stop characteristics and high heat input weld heat affected zone toughness Download PDF

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JPWO2009072559A1
JPWO2009072559A1 JP2009507634A JP2009507634A JPWO2009072559A1 JP WO2009072559 A1 JPWO2009072559 A1 JP WO2009072559A1 JP 2009507634 A JP2009507634 A JP 2009507634A JP 2009507634 A JP2009507634 A JP 2009507634A JP WO2009072559 A1 JPWO2009072559 A1 JP WO2009072559A1
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JP4612735B2 (en
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児島　明彦
明彦 児島
田中　洋一
洋一 田中
白幡　浩幸
浩幸 白幡
中島　清孝
清孝 中島
長井　嘉秀
嘉秀 長井
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Nippon Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/11—Making amorphous alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/50—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for welded joints

Abstract

この厚手高強度鋼板の製造方法は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含む連続鋳造スラブをＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱し、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を連続鋳造スラブに行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とし、次いで、加速冷却を適用して圧延原板を５００℃以下まで冷却して鋼板とし、前記連続鋳造スラブは、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たす。The manufacturing method of this thick high-strength steel sheet is mass%, C: 0.05-0.12%, Si: 0.3% or less, Mn: 1-2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al: 0.001 to 0.1%, Ti: 0.005 to 0.02%, After cooling a continuous cast slab containing N: 0.002 to 0.01% and O: 0.004% or less and containing iron and inevitable impurities as the balance to Ar3-200 ° C or less, 950 to 1100 ° C Next, rough rolling with a cumulative reduction amount of 30% or more at 900 ° C. or higher is performed on the continuous cast slab, and then finish rolling with a cumulative reduction amount of 50% or more at 700 ° C. or more is finished rolling. Both the start temperature and the finish rolling finish temperature are represented by the following formula {−0.5 × (slab heating Degree (° C.)) + 1325} (° C.) to obtain a rolled original sheet, and then applying accelerated cooling to cool the rolled original sheet to 500 ° C. or lower to obtain a steel sheet. The cast slab has a calculated value of B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation is 0% or less, and the carbon equivalent Ceq is in the range of 0.32 to 0.42%. Fulfill.

Description

本発明は、脆性破壊伝播停止特性と大入熱溶接での熱影響部（Ｈｅａｔ Ａｆｆｅｃｔｅｄ Ｚｏｎｅ：以下、ＨＡＺと称することがある）の靭性に優れた厚手高強度鋼板の製造方法、及び、脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板に関する。本発明に係る厚手高強度鋼板は、大型コンテナ船等の船舶向けとして主に使用されるが、建築、橋梁、タンク及び海洋構造物等、その他の溶接構造物に使用することも可能である。
本願は、２００７年１２月６日に出願された日本国特許出願第２００７−３１５８４０号に対し優先権を主張し、その内容をここに援用する。The present invention relates to a method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and toughness of a heat-affected zone (hereinafter sometimes referred to as HAZ) in high heat input welding, and brittle fracture The present invention relates to a thick high-strength steel sheet excellent in propagation stop characteristics and toughness of heat-affected zone with high heat input welding. The thick high-strength steel plate according to the present invention is mainly used for ships such as large container ships, but can also be used for other welded structures such as buildings, bridges, tanks and offshore structures.
This application claims priority to Japanese Patent Application No. 2007-315840 filed on Dec. 6, 2007, the contents of which are incorporated herein by reference.

船舶に代表される溶接構造物の近年のニーズとして、構造物の大型化、破壊に対する高い安全性、建造における溶接の高能率化、及び素材である鋼材の経済性等が挙げられる。このような動向を受け、溶接構造物に使用される鋼板に対して、（１）大きな板厚での高い強度、（２）良好な脆性破壊伝播停止特性、（３）良好な大入熱溶接ＨＡＺ靭性、及び（４）低い製造コスト等のニーズが高まりつつある。このため、大型コンテナ船などでは、降伏強度３９０ＭＰａ級（引張り強さ５１０ＭＰａ級）または４６０ＭＰａ級（引張強さ５７０ＭＰａ級）の船体構造用鋼板などが使用されるようになってきている。
具体的には、非特許文献１などのように、大型コンテナ船などの大型船舶に用いられる鋼板に対して、（１）板厚５０〜８０ｍｍの厚手鋼板（以下、厚手材と称することがある）での降伏強度３９０〜４６０ＭＰａ級（すなわち引張強度５１０〜５７０ＭＰａ級）の確保、（２）脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００（以下、アレスト性指標Ｔｋｃａ＝６０００と称することがある）≦−１０℃の確保、（３）溶接入熱量が２０ｋＪ／ｍｍ以上の溶接部のＨＡＺ靭性（シャルピー衝撃吸収エネルギー）ｖＥ（−２０℃）≧４７Ｊの確保、及び（４）高価合金元素の低減（Ｎｉ量≦１％等）を同時に満たすことが要求される。Recent needs for welded structures represented by ships include large-sized structures, high safety against destruction, high efficiency of welding in construction, and economical efficiency of steel materials. In response to these trends, (1) high strength with large plate thickness, (2) good brittle fracture propagation stop characteristics, and (3) good high heat input welding for steel plates used in welded structures There is a growing need for HAZ toughness and (4) low manufacturing costs. For this reason, in large container ships and the like, steel plates for hull structures having a yield strength of 390 MPa class (tensile strength of 510 MPa class) or 460 MPa class (tensile strength of 570 MPa class) have been used.
Specifically, as in Non-Patent Document 1, etc., (1) a thick steel plate having a thickness of 50 to 80 mm (hereinafter, referred to as a thick material) with respect to a steel plate used in a large vessel such as a large container ship. ) Yield strength of 390 to 460 MPa class (that is, tensile strength of 510 to 570 MPa class), (2) Temperature T kca = 6000 (hereinafter referred to as arrestability index) at which brittle fracture propagation stop characteristic Kca is 6000 N / mm 1.5 (Sometimes referred to as T kca = 6000 ) ≦ −10 ° C., (3) HAZ toughness (Charpy impact absorption energy) vE (−20 ° C.) ≧ 47 J of welded portion with a welding heat input of 20 kJ / mm or more And (4) It is required to satisfy simultaneously the reduction of expensive alloy elements (Ni amount ≦ 1%, etc.).

特許文献１は、船舶向け厚手高強度鋼板に関する技術の一例であり、この特許文献１には、板厚５０〜８０ｍｍを有しつつ、上記（１）、（３）及び（４）のニーズを部分的に満足できる技術が開示されている。しかしながら、特許文献１に記載の厚手高強度鋼板は、その実施例の記載からわかるように、上記（２）のニーズを満足できるような技術は示されていない。   Patent Document 1 is an example of a technology related to a thick high-strength steel sheet for ships. This Patent Document 1 has the needs of the above (1), (3), and (4) while having a plate thickness of 50 to 80 mm. A partially satisfactory technique is disclosed. However, the thick high-strength steel sheet described in Patent Document 1 does not show a technique that can satisfy the need (2), as can be seen from the description of the examples.

また、非特許文献２には、板厚が６５ｍｍと厚手の鋼板では、小型試験片によるシャルピー衝撃吸収エネルギーが、ｖＥ（−４０℃）＝１７０Ｊと十分に高くても、大型破壊試験で確認される脆性破壊伝播停止特性はＴｋｃａ＝６０００＝１８℃と不十分であることが示されている（同文献Ｆｉｇ．７参照）。これは、厚手鋼板では、小型試験片によるシャルピー衝撃吸収エネルギーｖＥ（−４０℃）を目安にして大型破壊試験で確認される脆性破壊伝播停止特性Ｔｋｃａ＝６０００≦−１０℃を保証することは困難であることを示している。すなわち、大型船舶向けの厚手高強度鋼板に要求される脆性破壊伝播停止特性を小型試験片によるシャルピー衝撃特性と関連付けて判定することは、従来の技術では困難であり、ＥＳＳＯ試験（ＷＥＳ ３００３準拠）に代表される全厚試験体の大型破壊試験を用いた方法でなければ、正確に評価することができなかった。In Non-Patent Document 2, in a thick steel plate with a thickness of 65 mm, even if Charpy impact absorption energy by a small test piece is sufficiently high as vE (−40 ° C.) = 170 J, it is confirmed by a large-scale fracture test. It has been shown that the brittle fracture propagation stop property is insufficient with T kca = 6000 = 18 ° C. (see FIG. 7 of the same document). This is because for thick steel plates, it is guaranteed that the brittle fracture propagation stop characteristic T kca = 6000 ≦ −10 ° C., which is confirmed by a large fracture test using Charpy impact absorption energy vE (−40 ° C.) by a small test piece as a guideline. It is difficult. In other words, it is difficult to determine the brittle fracture propagation stop characteristics required for thick high-strength steel sheets for large ships in association with the Charpy impact characteristics of small test pieces, and it is difficult with conventional technology, and an ESSO test (WES 3003 compliant) Unless it was a method using a large-scale destructive test of a full-thickness test body represented by (2), accurate evaluation could not be performed.

従来から、脆性破壊伝播停止特性は板厚に依存し、板厚が大きくなるほど当該特性が劣化することが知られていた。しかしながら、本発明が対象とするような５０ｍｍ以上の厚手材については、この板厚効果に関する実験データは皆無であり、厚手化に起因して脆性破壊伝播停止特性がどれくらい劣化するのかが不明であった。   Conventionally, it has been known that the brittle fracture propagation stop characteristics depend on the plate thickness, and that the properties deteriorate as the plate thickness increases. However, for thick materials of 50 mm or more as the object of the present invention, there is no experimental data on the plate thickness effect, and it is unclear how much the brittle fracture propagation stop characteristics deteriorate due to thickening. It was.

ところで、ＴＭＣＰ（Ｔｈｅｒｍｏ Ｍｅｃｈａｎｉｃａｌ Ｃｏｎｔｒｏｌ Ｐｒｏｃｅｓｓ）によって製造される厚手鋼板では、従来からボロン（Ｂ）添加による高強度化が図られてきた。Ｂの添加による効果としては、圧延後の加速冷却においてオーステナイト（γ）粒界に偏析した固溶Ｂが、変態時の焼入性を高めることが挙げられる。特許文献１では、ＢにＮｂを複合添加することによって高強度化を図っている。特許文献１の実施例に示されているように、この場合の圧延終了温度は９３０〜１０００℃と高いことが特徴であり、再結晶オーステナイト（再結晶γ）から加速冷却することを必須条件として、ＮｂとＢの複合効果を発揮させて高い焼入性を引き出すことにより、強度を高めている。一方、特許文献１では、圧延終了温度を９３０℃よりも低い未再結晶域として低温圧延を行った場合、靭性は満足するものの強度特性は満足できず、Ｎｂ−Ｂ複合効果による高強度化が難しいことも示されている。   By the way, in the thick steel plate manufactured by TMCP (Thermo Mechanical Control Process), high strength has been conventionally achieved by adding boron (B). As an effect by addition of B, solid solution B segregated at the austenite (γ) grain boundary in accelerated cooling after rolling increases the hardenability at the time of transformation. In Patent Document 1, an increase in strength is achieved by adding Nb to B in combination. As shown in the Examples of Patent Document 1, the rolling end temperature in this case is characterized by being as high as 930 to 1000 ° C., and accelerated cooling from recrystallized austenite (recrystallized γ) is an essential condition. , Nb and B are combined to bring out high hardenability by increasing the strength. On the other hand, in Patent Document 1, when low temperature rolling is performed with an unrecrystallized region having a rolling end temperature lower than 930 ° C., the toughness is satisfied but the strength characteristics cannot be satisfied, and the high strength due to the Nb—B composite effect is increased. It has also been shown to be difficult.

また、特許文献１では、大入熱溶接ＨＡＺにおけるＢ利用技術を開示しており、０．３０〜０．３８％のＣｅｑのもとで、γ中の固溶Ｂによる粒界フェライト抑制効果（焼入性向上効果）と、γ中のＢＮによる粒内フェライト促進効果（焼入性低減効果）を併用する有効性を示している。つまり、この場合、Ｂは焼入性に関して相反する二つの役割を担っている。以上から、特許文献１におけるＢ利用技術を要約すると、γ中の固溶Ｂによる焼入性向上効果を、直接焼入れ母材と大入熱溶接ＨＡＺで利用し、同時に、γ中の析出Ｂ（ここではＢＮ）による焼入性低減効果を、大入熱溶接ＨＡＺで利用している。   Moreover, in patent document 1, B utilization technique in the high heat input welding HAZ is disclosed. Under 0.30 to 0.38% of Ceq, the grain boundary ferrite suppression effect by solute B in γ ( It shows the effectiveness of combining the effect of improving hardenability and the effect of promoting intragranular ferrite (hardenability reducing effect) by BN in γ. That is, in this case, B plays two conflicting roles with respect to hardenability. From the above, the B utilization technology in Patent Document 1 is summarized as follows. The effect of improving the hardenability due to the solid solution B in γ is used in the directly quenched base material and the high heat input welding HAZ, and at the same time, the precipitation B in γ ( Here, the effect of reducing hardenability by BN) is used in high heat input welding HAZ.

また、本発明者等は、大入熱溶接ＨＡＺ靭性を高めるために、ＨＡＺの冷却過程でγ中に析出するＶＮをピン止め粒子（酸化物、硫化物）に複合析出させ、このＶＮ複合粒子がフェライト変態核として作用してＨＡＺ組織を微細化する発明を完成させ、特許文献２、３に開示している。また、非特許文献３に示されるように、Ｖ添加によって母材の強度が上昇する効果は広く知られている。
以上説明したように、ＢあるいはＶの添加によって、母材の強度が向上する効果と、大入熱溶接ＨＡＺの靭性が向上する効果が知られている。In addition, in order to increase the high heat input welding HAZ toughness, the inventors of the present invention caused VN precipitated in γ during the cooling process of HAZ to be compounded into pinning particles (oxide, sulfide), and this VN composite particle. Has completed the invention to refine the HAZ structure by acting as a ferrite transformation nucleus and disclosed in Patent Documents 2 and 3. As shown in Non-Patent Document 3, the effect of increasing the strength of the base material by adding V is widely known.
As described above, the effects of improving the strength of the base metal and the toughness of the high heat input welding HAZ by adding B or V are known.

一般に、母材やＨＡＺの靭性を高める希少な元素としてＮｉが知られており、上記（２）や（３）の観点からＮｉの有効利用が考えられる。しかしながら、Ｎｉは非常に高価な元素であり、その価格は近年著しく上昇している。また、Ｎｉを添加した鋼は表面疵が生じやすいため、その手入工程が発生するという問題がある。従って、Ｎｉ添加に関して、上記（４）のニーズと上記（２）及び（３）のニーズとの間で、その利害が対立する。また、上記（１）の観点から合金添加量を増加すると、炭素当量（Ｃｅｑ）が高まって大入熱溶接の場合のＨＡＺが硬化して脆化するので、上記（１）のニーズと上記（３）のニーズとの間で利害が対立する。さらに、上記（２）の観点からＴＭＣＰにおける変態前γ組織の微細化を追求すると、焼入性が低下して強度が減少するので、上記（１）のニーズと上記（２）のニーズとの間で利害が対立する。
このため、上述のような互いに利害が対立する上記（１）〜（４）の四つのニーズを同時に満足する鋼板の開発が強く求められていた。
特許第３５９９５５６号公報 特開２００５−２９８９００号公報 特開２００７−２６２５０８号公報 財団法人 日本海事協会「大型コンテナ船のＹＰ４７鋼の使用に関するガイドライン（２００８年１０月） 日本船舶海洋工学講演会論文集、２００６Ａ−Ｇ４−１０ ＣＡＭＰ−ＩＳＩＪ、６（１９９３）、ｐ６８４ In general, Ni is known as a rare element that enhances the toughness of a base material and HAZ. From the viewpoints of (2) and (3) above, Ni can be effectively used. However, Ni is a very expensive element, and its price has increased significantly in recent years. Further, since steel with Ni added tends to cause surface flaws, there is a problem that a care process is required. Therefore, regarding the addition of Ni, there is a conflict between the needs of (4) and the needs of (2) and (3). Further, if the amount of alloy addition is increased from the viewpoint of (1) above, the carbon equivalent (Ceq) increases and the HAZ in the case of high heat input welding hardens and becomes brittle. There is a conflict of interest with the needs of 3). Further, pursuing refinement of the pre-transformation γ structure in TMCP from the viewpoint of (2) above, the hardenability decreases and the strength decreases, so the needs of (1) and (2) above Conflicts of interest.
For this reason, there has been a strong demand for the development of a steel sheet that simultaneously satisfies the four needs (1) to (4), which have conflicting interests as described above.
Japanese Patent No. 3599556 JP 2005-298900 A JP 2007-262508 A Japan Maritime Association “Guidelines for Using YP47 Steel in Large Container Ships (October 2008) Proceedings of the Japan Marine Engineering Lecture, 2006A-G4-10 CAMP-ISIJ, 6 (1993), p684

本発明は上記問題に鑑みてなされたものであり、（１）板厚５０〜８０ｍｍ、降伏強度３９０〜４６０ＭＰａ級、かつ引張強度５１０〜５７０ＭＰａ級の厚手高強度で、（２）アレスト性指標Ｔｋｃａ＝６０００≦−１０℃の良好な脆性破壊伝播停止特性を有し、（３）溶接入熱量≧２０ｋＪ／ｍｍでもｖＥ（−２０℃）≧４７Ｊとなる良好な大入熱溶接ＨＡＺ靭性を有し、（４）高価合金元素の低減（Ｎｉ≦１％等）等による低い製造コストを実現できる、脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法、及び脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板を提供することを目的とする。The present invention has been made in view of the above problems. (1) Thickness and high strength of a plate thickness of 50 to 80 mm, a yield strength of 390 to 460 MPa, and a tensile strength of 510 to 570 MPa, and (2) an arrestability index T It has good brittle fracture propagation stop characteristics of kca = 6000 ≤ -10 ° C, and (3) has good large heat input weld HAZ toughness with vE (-20 ° C) ≥ 47J even when the welding heat input ≥ 20 kJ / mm (4) A method for producing a thick, high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness, which can realize low production cost due to reduction of expensive alloy elements (Ni ≦ 1%, etc.) It is another object of the present invention to provide a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness.

上記問題を解決するための本発明の要旨は以下のとおりである。
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含む連続鋳造スラブをＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱し、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を前記連続鋳造スラブに行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とし、次いで、加速冷却を適用して前記圧延原板を５００℃以下まで冷却して鋼板とする。前記連続鋳造スラブは、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たす。
ここで、強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量ＯＴｉ（％）を、下記式（１）で表される量としたとき、有効Ｂ量：Ｂｅｆ（％）は、下記式（２）で表される。また炭素当量Ｃｅｑ（％）は、下記式（３）で表され、Ａｒ３は、下記式（４）で表される。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（１）｛但し、式（１）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）｛但し、式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（３）
Ａｒ３（℃）＝（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ） ・・・（４）
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法では、前記加速冷却の後、さらに、３５０〜７００℃で５〜６０分の焼戻し熱処理を施してもよい。
前記連続鋳造スラブの前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、前記連続鋳造スラブは、さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有してもよい。
前記連続鋳造スラブは、さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有してもよい。
前記連続鋳造スラブは、さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有してもよい。The gist of the present invention for solving the above problems is as follows.
The manufacturing method of the thick high-strength steel sheet excellent in the brittle fracture propagation stopping characteristics and the high heat input welding heat-affected zone toughness of the present invention is mass%, C: 0.05 to 0.12%, Si: 0.3%. Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al: 0 0.001 to 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, and O: 0.004% or less, with iron and inevitable impurities as the balance The continuous cast slab containing is cooled to Ar 3 −200 ° C. or lower, then reheated to 950 to 1100 ° C., and then subjected to rough rolling with a cumulative reduction of 30% or higher at 900 ° C. or higher, Next, finish rolling with a cumulative reduction amount of 50% or more at 700 ° C. or higher is determined as the finish rolling start temperature. The finish rolling finish temperature is the same as the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.). The rolling original sheet is cooled to 500 ° C. or less by applying cooling to obtain a steel sheet. In the continuous cast slab, the calculated amount of B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation is 0% or less, and the carbon equivalent Ceq is 0.32 to 0.42%. Fill the range.
Here, when the residual oxygen amount O Ti (%) that remains after deoxidation by the strong deoxidation element and can be deoxidized by Ti, which is the weak deoxidation element, is effective when the amount is represented by the following formula (1) B amount: Bef (%) is represented by the following formula (2). The carbon equivalent Ceq (%) is represented by the following formula (3), and Ar 3 is represented by the following formula (4).
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1) {However, in the formula (1), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (2) {However, in the formula (2), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (3)
Ar 3 (° C.) = (910-310C-80Mn-20Cu-55Ni-80Mo) (4)
In the method for producing a thick high-strength steel sheet having excellent brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention, after the accelerated cooling, tempering heat treatment is further performed at 350 to 700 ° C. for 5 to 60 minutes. You may give it.
The content of S in the continuous cast slab is 0.0005 to 0.005%, and the content of O is 0.001 to 0.004%. Thus, one or two of Ca: 0.0003 to 0.004% and Mg: 0.0003 to 0.004% may be contained.
The continuous cast slab further comprises, in mass%, Ni: 0.01 to 1%, Cu: 0.01 to 1%, Cr: 0.01 to 1%, Mo: 0.01 to 0.5%, And Nb: You may contain 1 type (s) or 2 or more types in 0.003-0.03%.
The continuous cast slab may further contain one or two of REM: 0.0003 to 0.02% and Zr: 0.0003 to 0.02% by mass%.

本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含み、強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量を、下記式（５）で表される量としたとき、下記式（６）で表される、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、さらに、下記式（７）で表される炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たし、板厚が５０〜８０ｍｍであり、降伏強度が３９０〜４６０ＭＰａ級であり、脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００が−１０℃以下であり、溶接入熱量が２０ｋＪ／ｍｍ以上の大入熱溶接部のＨＡＺ靭性の指標であるシャルピー衝撃吸収エネルギーｖＥ（−２０℃）が４７Ｊ以上である。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（５）｛但し、式（５）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（６）｛但し、式（６）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（７）
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板では、前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有してもよい。
さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有してもよい。
さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有してもよい。The thick high-strength steel sheet having excellent brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention is mass%, C: 0.05 to 0.12%, Si: 0.3% or less, Mn : 1-2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003-0.003%, V: 0.01-0.15%, Al: 0.001- 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, and O: 0.004% or less, including iron and inevitable impurities as the balance, strong When the amount of residual oxygen remaining after deoxidation by the deoxidation element and deoxidized by Ti, which is a weak deoxidation element, is represented by the following formula (5), it is represented by the following formula (6). The calculated value of B amount {effective B amount: Bef (%)} that dissolves in the austenite substrate before transformation is 0% or less. Furthermore, the following formula (7) Carbon equivalent Ceq represented satisfies the range of 0.32 to 0.42%, a plate thickness of 50 to 80 mm, the yield strength is 390~460MPa grade, brittle fracture propagation stop characteristics Kca is 6000 N / mm 1 The temperature T kca = 6000 to be 0.5 is −10 ° C. or less, and the Charpy impact absorption energy vE (−20 ° C.), which is an index of the HAZ toughness of a high heat input weld having a welding heat input of 20 kJ / mm or more, is 47 J. That's it.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (5) {However, in the formula (5), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (6) {However, in the formula (6), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (7)
In the thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness of the present invention, the content of S is 0.0005 to 0.005%, and the content of O is 0.001 to 0.004%, and further, by mass%, containing one or two of Ca: 0.0003 to 0.004% and Mg: 0.0003 to 0.004% Also good.
Furthermore, by mass%, Ni: 0.01-1%, Cu: 0.01-1%, Cr: 0.01-1%, Mo: 0.01-0.5%, and Nb: 0.003 You may contain 1 type, or 2 or more types in -0.03%.
Furthermore, you may contain 1 type or 2 types in REM: 0.0003-0.02% and Zr: 0.0003-0.02% by the mass%.

本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法、及び脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板によれば、（１）板厚５０〜８０ｍｍ、降伏強度３９０〜４６０ＭＰａ級（すなわち引張強度５１０〜５７０ＭＰａ級）の厚手高強度で、（２）アレスト性指標Ｔｋｃａ＝６０００≦−１０℃の良好な脆性破壊伝播停止特性を有し、（３）溶接入熱量≧２０ｋＪ／ｍｍでもｖＥ（−２０℃）≧４７Ｊとなる良好な大入熱溶接ＨＡＺ靭性を有し、（４）高価合金元素の低減（Ｎｉ≦１％等）等による低い製造コストを実現できる。
このような本発明による厚手高強度鋼板が大型船舶をはじめとする各種の溶接構造物に使用されることで、溶接構造物の大型化、破壊に対する高い安全性、建造における溶接の高能率化、及び素材である鋼材の経済性等々が同時に満たされることから、その産業上の効果は計り知れない。Manufacturing method of thick high-strength steel sheet excellent in brittle fracture propagation stop characteristics and large heat input welding heat-affected zone toughness, and thick high-strength steel sheet excellent in brittle fracture propagation stop characteristics and large heat input welding heat-affected zone toughness (1) Thickness and high strength of 50 to 80 mm in thickness and yield strength of 390 to 460 MPa (that is, tensile strength of 510 to 570 MPa), and (2) good arrestability index T kca = 6000 ≦ −10 ° C. (3) Good high heat input welding HAZ toughness with vE (−20 ° C.) ≧ 47 J even when the welding heat input ≧ 20 kJ / mm, (4) Low manufacturing costs due to reduction (Ni ≦ 1%, etc.) can be realized.
Such a thick high-strength steel sheet according to the present invention is used for various welded structures including large ships, so that the welded structures are enlarged, high safety against breakage, high efficiency of welding in construction, In addition, since the economics of the steel material, etc. are satisfied at the same time, the industrial effects are immeasurable.

以下、本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法、及び脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の実施の形態について説明する。
なお、この実施形態は、発明の趣旨をより良く理解させるために詳細に説明するものであるから、特に指定の無い限り、本発明を限定するものではない。Hereinafter, the manufacturing method of a thick high strength steel sheet excellent in brittle fracture propagation stop characteristics and large heat input welding heat affected zone toughness of the present invention, and thick high excellent in brittle fracture propagation stop characteristics and large heat input welding heat affected zone toughness. An embodiment of a strength steel plate will be described.
In addition, since this embodiment is described in detail for better understanding of the gist of the invention, the present invention is not limited unless otherwise specified.

＜鋼板製造条件（製造方法）＞
船舶等の溶接構造物に使用される鋼板においては、（１）大きな板厚での高い強度、（２）良好な脆性破壊伝播停止特性、（３）良好な大入熱溶接ＨＡＺ靭性、及び（４）低い製造コスト等のニーズが高まっている。
このようなニーズに対し、本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含む連続鋳造スラブを、連続鋳造後にＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱する工程と、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を連続鋳造スラブに行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とする工程と、次いで、加速冷却を適用して圧延原板を５００℃以下まで冷却して鋼板とする工程とを有する。前記連続鋳造スラブは、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たす。
ここで、強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量ＯＴｉ（％）を、下記式（１）で表される量としたとき、有効Ｂ量：Ｂｅｆ（％）は、下記式（２）で表される。また炭素当量Ｃｅｑ（％）は、下記式（３）で表され、Ａｒ３は、下記式（４）で表される。
また、スラブ加熱温度とは、連続鋳造スラブを再加熱する際の温度（再加熱温度）である。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（１）｛但し、式（１）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）｛但し、式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（３）
Ａｒ３（℃）＝（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ） ・・・（４）
なお、本明細書において、式中の元素記号は、連続鋳造スラブ又は厚手高強度鋼板中のその元素の含有量（質量％）を示す。
また、本発明において連続鋳造スラブの製造方法は特に限定されない。例えば、高炉、転炉や電炉等による溶製に引き続き、各種の２次精練で目的の成分含有量になるように成分調整を行い、次いで通常の連続鋳造によって製造される。<Steel plate manufacturing conditions (manufacturing method)>
In steel plates used for welded structures such as ships, (1) high strength at large plate thickness, (2) good brittle fracture propagation stop properties, (3) good high heat input weld HAZ toughness, and ( 4) Needs such as low manufacturing costs are increasing.
For such needs, the method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention is in mass%, and C: 0.05 to 0.12. %, Si: 0.3% or less, Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al: 0.001-0.1%, Ti: 0.005-0.02%, N: 0.002-0.01%, and O: 0.004% or less, A continuous cast slab containing iron and inevitable impurities as the balance is cooled to Ar 3 −200 ° C. or lower after continuous casting and then reheated to 950 to 1100 ° C., and then the cumulative reduction amount is 900 ° C. or higher and the cumulative reduction amount is 30 % Or more rough rolling is performed on the continuous cast slab, and then the cumulative reduction at 700 ° C or higher The finish rolling which is 50% or more, and the finish rolling start temperature and the finish rolling end temperature are both below the temperature represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.) And a step of forming a rolled original sheet under the above conditions, and then a step of applying accelerated cooling to cool the rolled original sheet to 500 ° C. or less to obtain a steel sheet. In the continuous cast slab, the calculated amount of B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation is 0% or less, and the carbon equivalent Ceq is 0.32 to 0.42%. Fill the range.
Here, when the residual oxygen amount O Ti (%) that remains after deoxidation by the strong deoxidation element and can be deoxidized by Ti, which is the weak deoxidation element, is effective when the amount is represented by the following formula (1) B amount: Bef (%) is represented by the following formula (2). The carbon equivalent Ceq (%) is represented by the following formula (3), and Ar 3 is represented by the following formula (4).
The slab heating temperature is a temperature (reheating temperature) when reheating the continuously cast slab.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1) {However, in the formula (1), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (2) {However, in the formula (2), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (3)
Ar 3 (° C.) = (910-310C-80Mn-20Cu-55Ni-80Mo) (4)
In addition, in this specification, the element symbol in a type | formula shows content (mass%) of the element in a continuous casting slab or a thick high-strength steel plate.
In the present invention, the method for producing the continuous cast slab is not particularly limited. For example, subsequent to melting by a blast furnace, a converter, an electric furnace, etc., the components are adjusted so that the desired component content is obtained by various secondary scouring, and then manufactured by ordinary continuous casting.

本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法では、上記化学成分組成において、上記各元素のうち、Ｓの含有量の下限を０．０００５％、Ｏの含有量の下限を０．００１％とすることができる。さらに、必要に応じて、Ｃａ：０．０００３〜０．００４％、Ｍｇ：０．０００３〜０．００４％、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、Ｎｂ：０．００３〜０．０３％、ＲＥＭ：０．０００３〜０．０２％、及びＺｒ：０．０００３〜０．０２％のうちの１種または２種以上を選択的に含有することができる。
なお、ＲＥＭとは、希土類金属であり、Sc，Y，及びランタノイドのLa，Ce，Pr，Nd，Pm，Sm，Eu，Gd，Tb，Dy，Ho，Er，Tm，Yb，及びLuから選択される１種以上である。In the method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention, the lower limit of the content of S among the above elements in the chemical component composition is 0. .0005%, and the lower limit of the O content can be 0.001%. Furthermore, if necessary, Ca: 0.0003 to 0.004%, Mg: 0.0003 to 0.004%, Ni: 0.01 to 1%, Cu: 0.01 to 1%, Cr: 0 0.01-1%, Mo: 0.01-0.5%, Nb: 0.003-0.03%, REM: 0.0003-0.02%, and Zr: 0.0003-0.02% 1 type or 2 types or more can be selectively contained.
REM is a rare earth metal selected from Sc, Y, and lanthanoids such as La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu. 1 or more types.

本発明の要点は、ＴＭＣＰ型で製造する厚手鋼板において、強度、脆性破壊伝播停止特性、大入熱溶接ＨＡＺ靭性、及び低い製造コスト等を同時に満足するために、ＢとＶを複合添加することを特徴とし、これら窒化物形成元素と結合するＮを精緻に制御することでオーステナイト（γ）中のＢとＶの存在状態を最適化し、母材と大入熱溶接ＨＡＺの変態組織を制御する技術である。
具体的には、γ中のＢの存在状態に関しては、母材と大入熱溶接ＨＡＺの両方において、固溶Ｂを存在させずに全てのＢをＢＮとして析出させる技術思想である。γ中のＶの存在状態に関しては、母材では固溶Ｖとして利用し、大入熱溶接ＨＡＺでは析出Ｖ（ＶＮ等）として利用する技術思想である。The main point of the present invention is to add B and V together in order to simultaneously satisfy the strength, brittle fracture propagation stop characteristics, high heat input welding HAZ toughness, low production cost, etc. By optimizing N bonding with these nitride-forming elements, the state of B and V in austenite (γ) is optimized, and the transformation structure of the base metal and high heat input weld HAZ is controlled. Technology.
Specifically, regarding the existence state of B in γ, it is a technical idea that all B precipitates as BN without the presence of solid solution B in both the base metal and the high heat input welding HAZ. Regarding the existence state of V in γ, it is a technical idea that it is used as a solid solution V in the base metal and as a precipitation V (VN or the like) in the high heat input welding HAZ.

以下、詳細を説明する。
まず、本発明における最大の技術課題である脆性破壊伝播停止特性を満足するため、厚手鋼板の結晶粒径を極限まで微細化するＴＭＣＰ条件を検討した。
ここで、脆性破壊が結晶学的に同一の結晶面（へき開面：体心立方構造の鉄では｛１００｝面に対応）で生じる最小単位は破面単位と呼ばれ、この破面単位に対応するサイズの金属組織単位を本発明では「結晶粒径」と呼ぶこととする。
ＴＭＣＰにおける低温加熱と低温圧延を徹底して変態前γの微細化を限界まで追求すれば、板厚が５０〜８０ｍｍである厚手鋼板であっても結晶粒径が充分に微細化し、脆性破壊伝播停止特性が目標を満足できることが明らかとなった。その条件は、Ａｒ３（℃）が次式（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ）で計算されるとき、連続鋳造スラブを｛Ａｒ３（℃）−２００（℃）｝以下の温度まで冷却した後に１１００℃以下に低温加熱（再加熱）し、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度（℃）及び仕上圧延終了温度（℃）が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行い、次いで、加速冷却を適用して５００℃以下まで冷却することである。Details will be described below.
First, in order to satisfy the brittle fracture propagation stop characteristic which is the greatest technical problem in the present invention, the TMCP conditions for minimizing the crystal grain size of the thick steel plate were studied.
Here, the smallest unit in which brittle fracture occurs on the same crystallographic crystal plane (cleavage plane: corresponding to the {100} plane in iron with a body-centered cubic structure) is called the fracture surface unit, and corresponds to this fracture surface unit. In the present invention, the metal structure unit of the size is referred to as “crystal grain size”.
If thorough low temperature heating and rolling at TMCP are pursued and the refinement of γ before transformation is pursued to the limit, the grain size will be sufficiently refined even for thick steel plates with a thickness of 50 to 80 mm, and brittle fracture propagation will occur. It became clear that the stopping characteristics can meet the target. The condition is that when Ar 3 (° C.) is calculated by the following formula (910-310C-80Mn-20Cu-55Ni-80Mo), the temperature of the continuously cast slab is {Ar 3 (° C.) − 200 (° C.)} or less. After cooling to 1100 ° C. or lower, the steel sheet is cooled to 1100 ° C. or lower (reheated), then, rough rolling is performed at 900 ° C. or higher and the cumulative reduction amount is 30% or higher. The temperature at which the finish rolling start temperature (° C.) and the finish rolling end temperature (° C.) are both represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.) It is performed under the following conditions, and then cooled to 500 ° C. or lower by applying accelerated cooling.

低温加熱、低温圧延を徹底するＴＭＣＰの第一条件として、連続鋳造後のスラブ（連続鋳造スラブ）をＡｒ３−２００℃以下に冷却してγ（オーステナイト）→α（フェライト）変態させ、その後に１１００℃以下に低温加熱（再加熱）することでα→γ変態させる。この製造条件を適用する理由は、加熱時のγを徹底的に整細粒化するためである。
スラブを、｛Ａｒ３（℃）−２００（℃）｝を超える高温から再加熱すると、スラブ内部でγ→α変態が未完了のうちに再加熱されて鋳造時の粗大γが残存してしまう。上記式（４）はスラブが連続鋳造されて冷却する際の極めて小さな冷却速度について成り立つ関係であり、厚板圧延のように冷却速度が相対的に大きい場合には適用されない。
スラブの再加熱温度が１１００℃を超えるような高温加熱を行うと、ＴｉＮのオストワルド成長が始まるため、ピン止め効果が低減して整細粒γを安定的に確保することが難しくなる。加熱時のγを徹底的に整細粒化できなければ、現実的なスラブ厚みの制約下（通常は２００〜４００ｍｍ）において、圧延条件を如何に工夫したとしても、板厚が５０〜８０ｍｍである鋼板の変態前γを十分に微細化することは困難である。As the first condition of TMCP for thorough low-temperature heating and low-temperature rolling, the slab after continuous casting (continuous casting slab) is cooled to below Ar 3 -200 ° C. to transform γ (austenite) → α (ferrite), and then The α → γ transformation is performed by low-temperature heating (reheating) to 1100 ° C. or lower. The reason for applying this manufacturing condition is to thoroughly refine γ during heating.
When the slab is reheated from a high temperature exceeding {Ar 3 (° C.) − 200 (° C.)}, the γ → α transformation is reheated incomplete in the slab, and coarse γ during casting remains. . The above formula (4) is a relationship that holds for a very small cooling rate when the slab is continuously cast and cooled, and is not applied when the cooling rate is relatively high, such as thick plate rolling.
When high-temperature heating is performed such that the reheating temperature of the slab exceeds 1100 ° C., Ostwald growth of TiN starts, so that the pinning effect is reduced and it becomes difficult to stably secure the fine grain γ. If γ during heating cannot be thoroughly refined, the plate thickness is 50 to 80 mm no matter how the rolling conditions are devised under the realistic slab thickness constraints (usually 200 to 400 mm). It is difficult to sufficiently refine γ before transformation of a certain steel sheet.

低温加熱、低温圧延を徹底するＴＭＣＰの第二条件として、９００℃以上で累積圧下量が３０％以上である粗圧延を行う。この製造条件を適用する理由は、再結晶域での圧延によって加熱時よりもさらに整細粒なγを得るためである。
粗圧延が９００℃未満であったり、また、累積圧下量が３０％未満であると、再結晶が不十分となって歪誘起粒成長が起こり、加熱時の初期γよりもむしろ粗大になる恐れがある。As a second condition of TMCP for thorough low-temperature heating and low-temperature rolling, rough rolling is performed at 900 ° C. or higher and the cumulative reduction amount is 30% or higher. The reason for applying this production condition is to obtain γ that is finer than that during heating by rolling in the recrystallization region.
If rough rolling is less than 900 ° C. or if the cumulative reduction is less than 30%, recrystallization is insufficient and strain-induced grain growth occurs, which may become coarse rather than initial γ during heating. There is.

低温加熱、低温圧延を徹底するＴＭＣＰの第三条件として、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度（℃）及び仕上圧延終了温度（℃）を、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行う。この製造条件を適用する理由は、粗圧延で十分に整細粒化した再結晶粒を未再結晶域圧延することで、γ粒を延伸化させて粒界の面積を増やすとともに粒界を活性化させ、さらにγ粒内に変形帯を導入し、変態前γにおける核生成サイト密度と核生成頻度を限界まで高めるためである。
仕上圧延の累積圧下量が５０％未満であったり、また、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下の条件を満たさない場合は、変態前γの微細化が不十分となる。
上記式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下の条件の金属学的な意味としては、加熱温度が高く初期γが粗大であるほど、仕上圧延をより低温で行って未再結晶域圧延を強化する必要があることを示している。例えば、スラブ加熱温度が１１００℃ならば仕上圧延を７７５℃以下で行う必要があり、スラブ加熱温度が１０００℃ならば８２５℃以下で圧延を行う必要がある。このように、スラブ加熱温度に連動させて仕上圧延温度を規制する極めて厳格なＴＭＣＰ条件を適用しないと、厚手鋼板で良好な脆性破壊伝播停止特性を安定して確保することはできない。
７００℃よりも低温域で仕上圧延を行うと、圧延中あるいは加速冷却までの待機時間中に鋼板の表層側が変態を開始してしまい、表層部組織が軟化すると同時に粗大化してしまい、強度と脆性破壊伝播停止特性が劣化する。As a third condition of TMCP for thorough low-temperature heating and low-temperature rolling, finish rolling with a cumulative reduction amount of 50% or more at 700 ° C or higher, both finish rolling start temperature (° C) and finish rolling end temperature (° C), , Under the condition of not more than the temperature represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.). The reason for applying this manufacturing condition is that the recrystallized grains that have been sufficiently refined by rough rolling are rolled in the non-recrystallized zone, thereby extending the γ grains to increase the area of the grain boundaries and activating the grain boundaries. In order to increase the density of nucleation sites and the frequency of nucleation in the γ before transformation to the limit.
When the cumulative reduction amount of finish rolling is less than 50% or does not satisfy the conditions below the temperature represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.) Further, the refinement of γ before transformation becomes insufficient.
As the metallurgical meaning of the condition below the temperature represented by the above formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.), the heating temperature is high and the initial γ is coarser. It shows that finish rolling needs to be performed at a lower temperature to strengthen non-recrystallized zone rolling. For example, if the slab heating temperature is 1100 ° C., it is necessary to perform finish rolling at 775 ° C. or less, and if the slab heating temperature is 1000 ° C., it is necessary to perform rolling at 825 ° C. or less. Thus, unless the extremely strict TMCP condition that regulates the finish rolling temperature in conjunction with the slab heating temperature is applied, it is impossible to stably ensure good brittle fracture propagation stop characteristics with the thick steel plate.
When finish rolling is performed at a temperature lower than 700 ° C., the surface layer side of the steel sheet starts transformation during rolling or waiting time until accelerated cooling, and the surface layer structure is softened and coarsened at the same time, resulting in strength and brittleness. Destruction propagation stop characteristics deteriorate.

低温加熱、低温圧延を徹底するＴＭＣＰの第四条件として、加速冷却を適用して５００℃以下まで冷却する。この製造条件を適用する理由は、上述のごとく加熱、圧延条件を徹底して変態前γを限界まで微細化しても、その後の冷却が空冷であるとγ→α変態時の過冷度が小さく、結晶粒径が十分に微細化できないからである。
加速冷却を５００℃よりも高温で停止すると、板厚表層に比べて温度の高い板厚内部では、変態の途中で加速冷却が終了して空冷になるため、板厚内部の結晶粒径が十分に微細化できない。As the fourth condition of TMCP for thorough low temperature heating and low temperature rolling, accelerated cooling is applied to cool to 500 ° C. or lower. The reason for applying this manufacturing condition is that, as described above, even if the pre-transformation γ is refined to the limit through thorough heating and rolling conditions, if the subsequent cooling is air cooling, the degree of supercooling during the γ → α transformation is small. This is because the crystal grain size cannot be sufficiently reduced.
If accelerated cooling is stopped at a temperature higher than 500 ° C., the accelerated cooling is completed in the middle of the transformation and air cooling occurs inside the plate thickness, which is higher than the surface thickness of the plate thickness. Can not be refined.

以上が、低Ｎｉを前提に結晶粒径を十分に微細化して脆性破壊伝播停止特性を満足するためのＴＭＣＰ条件であり、上記（２）と（４）のニーズを満足できる。
しかしながら、上述のＴＭＣＰ条件では、変態前γの徹底した微細化と厚手鋼板特有の小さな冷却速度が相俟って、変態時の焼入性が大幅に低下する問題が生じる。その結果、ベイナイト／フェライト混合組織におけるベイナイト分率が減少してフェライト分率が増加し、所定の引張強度を確保することが難しい。同時に、このようなＴＭＣＰ条件ではγ中の固溶Ｂに起因する焼入性も不安定となり、強度不足に加えて強度ばらつきが大きくなる問題が判明した。このように、上述のＴＭＣＰ条件では、上記（１）のニーズを満たせないことが新たな課題として浮上した。
強度のばらつきの第一の理由は、後述する有効Ｂ量（Ｂｅｆ）で見積もられるγ中の固溶Ｂ量が、大量生産時の鋼成分変動（Ｏ量、強脱酸元素量、Ｔｉ量、Ｎ量、Ｂ量の変動）に起因して増減するためである。第二の理由は、低温圧延された未再結晶域γ状態では、圧延条件や圧延後の加速冷却開始までの待機時間に依存して、鉄炭硼化物（Ｆｅ２３（Ｃ，Ｂ）６等）の歪誘起析出量が変動し、その裏返しとしてγ中の固溶Ｂ量が増減するためである。以上のように、上述のＴＭＣＰ条件ではＢ焼入性に頼って母材強度を安定確保することは容易でなく、Ｂ焼入性以外の強化手段を利用する必要性が生じた。The above is the TMCP condition for sufficiently reducing the crystal grain size on the premise of low Ni and satisfying the brittle fracture propagation stop characteristics, and can satisfy the above-mentioned needs (2) and (4).
However, under the above-mentioned TMCP conditions, the thorough refinement of γ before transformation and the small cooling rate peculiar to thick steel plates cause a problem that the hardenability during transformation is significantly lowered. As a result, the bainite fraction in the bainite / ferrite mixed structure decreases, the ferrite fraction increases, and it is difficult to ensure a predetermined tensile strength. At the same time, it was found that the hardenability caused by the solid solution B in γ becomes unstable under such TMCP conditions, and the strength variation becomes large in addition to insufficient strength. As described above, it has emerged as a new problem that the above-mentioned TMCP condition cannot satisfy the above-mentioned need (1).
The first reason for the variation in strength is that the amount of solute B in γ estimated by the effective B amount (Bef), which will be described later, is a change in steel components during mass production (O amount, strong deoxidizing element amount, Ti amount, This is because it increases or decreases due to fluctuations in the N amount and the B amount. The second reason is that, in the non-recrystallized region γ state that has been cold-rolled, iron boride (Fe 23 (C, B) 6, etc.), depending on the rolling conditions and the waiting time until the start of accelerated cooling after rolling. This is because the amount of strain-induced precipitation of) fluctuates and the amount of dissolved B in γ increases or decreases as the reverse. As described above, under the above TMCP conditions, it is not easy to ensure a stable base metal strength by relying on B hardenability, and it is necessary to use strengthening means other than B hardenability.

そこで本発明では、上記（１）のニーズを満たすために、母材強度を安定かつ十分に確保することを狙って、下記の二つの手段を講じる。
第一の手段は、ＴＭＣＰにおいてγ中に固溶Ｂを存在させず、全てのＢをＢＮとして析出させることで、γ中の固溶Ｂ量の変動に起因する焼入性の不安定性を排除する。これは、従来のＢ利用技術とは全く反対の考え方であり、母材強度のためにＢ焼入性を使わない技術思想である。これにより、大量生産における強度ばらつきを抑えることができる。具体的には、後述する有効Ｂ量（Ｂｅｆ）を０％以下に制御する。本発明でＢを添加する意義は大入熱溶接ＨＡＺにあり、この点については後述する。
第二の手段は、Ｖ炭化物による析出強化を利用して母材強度を高める。
上述のＴＭＣＰ条件では、０．０１％のＶを添加することで板厚７０ｍｍ材の引張強度が１０ＭＰａ程度上昇することが判明し、Ｖ添加が極めて有効な強化手段であることが定量的に明らかとなった。これは、低温加熱と低温圧延を徹底して十分に微細化したベイナイト／フェライト混合組織が、加速冷却や焼き戻し処理においてＶ炭化物（ＶＣ、Ｖ４Ｃ３等）が微細高密度に析出する素地として好適であるためである。本発明でＶを添加するもう一つの意義は、大入熱溶接ＨＡＺにあり、この点については後述する。Therefore, in the present invention, in order to satisfy the above-mentioned need (1), the following two measures are taken with the aim of ensuring the base material strength stably and sufficiently.
The first measure eliminates hardenability instability caused by fluctuations in the amount of solid solution B in γ by precipitating all B as BN in TMCP without having solid solution B present in γ. To do. This is a concept opposite to the conventional B utilization technology, and is a technical idea that does not use B hardenability for the strength of the base material. Thereby, intensity variation in mass production can be suppressed. Specifically, the effective B amount (Bef) described later is controlled to 0% or less. The significance of adding B in the present invention is high heat input welding HAZ, which will be described later.
The second means increases the base material strength by utilizing precipitation strengthening by V carbide.
Under the above-mentioned TMCP conditions, it has been found that the addition of 0.01% V increases the tensile strength of the material having a thickness of 70 mm by about 10 MPa, and quantitatively reveals that the addition of V is an extremely effective strengthening means. It became. This is because the bainite / ferrite mixed structure, which has been sufficiently refined through thorough low-temperature heating and low-temperature rolling, is a substrate on which V carbides (VC, V 4 C 3 and the like) precipitate in a fine and high density during accelerated cooling and tempering treatment. It is because it is suitable as. Another significance of adding V in the present invention lies in high heat input welding HAZ, which will be described later.

以上説明したように、ＴＭＣＰにおいてＢ焼入性を使わずにＶ添加で母材強度を確保するために、Ｂを除く鋼成分の焼入性の目安として採用する炭素当量Ｃｅｑを０．３２％以上確保したうえで、有効Ｂ量Ｂｅｆを０％以下に制御し、Ｖを０．０１％以上添加し、加熱温度を９５０℃以上に制御し、加速冷却を５００℃以下まで行うことが必要である。
Ｃｅｑが０．３２％未満の場合、Ｖを添加したとしても母材強度を安定確保することが難しい。さらに、ＨＡＺ軟化が大きくなって溶接継手の引張強度が不足する恐れがある。
上記式（２）で算出される有効Ｂ量が０％を超えて多い数値の場合、γ中の固溶Ｂが存在してＢ焼入性が発現してしまい、強度がばらつく恐れがある。
加熱温度が９５０℃未満の場合、Ｖ炭窒化物の溶体化が不十分となり、γ中の固溶Ｖが不足することで、加速冷却や焼き戻し処理で析出するＶ炭化物が不足して母材強度を安定確保できない。
加速冷却ではなく空冷を適用すると、冷却速度が小さすぎてフェライトが粗大化すると同時にベイナイト分率が減少し、変態強化が十分に得られない。
加速冷却を５００℃よりも高温で停止すると、温度の高い板厚内部は変態途中で加速冷却が終了してしまうため、板厚内部の変態強化が十分に得られない。
加速冷却においては、０．３ｍ３／ｍ２／ｍｉｎ以上の水量密度を確保することが、強度と靭性を両立する微細なベイナイト／フェライト組織を得るのに好ましい。
以上が、脆性破壊発生特性を重視したＴＭＣＰ条件において、低Ｎｉを前提に強度を満足できる技術であり、これによって上記（１）、（２）、（４）のニーズを同時に満足することができる。As described above, in order to secure the base metal strength by adding V without using B hardenability in TMCP, the carbon equivalent Ceq adopted as a measure of hardenability of steel components excluding B is 0.32%. After securing the above, it is necessary to control the effective B amount Bef to 0% or less, to add V 0.01% or more, to control the heating temperature to 950 ° C or more, and to perform accelerated cooling to 500 ° C or less. is there.
When Ceq is less than 0.32%, it is difficult to stably secure the base material strength even if V is added. Furthermore, there is a possibility that the HAZ softening becomes large and the tensile strength of the welded joint is insufficient.
In the case where the effective B amount calculated by the above formula (2) is a large value exceeding 0%, solid solution B in γ exists and B hardenability is exhibited, which may cause variations in strength.
When the heating temperature is less than 950 ° C., the solution of V carbonitride is insufficient, and the solid solution V in γ is insufficient, so that the V carbide precipitated by accelerated cooling and tempering is insufficient, and the base material. The strength cannot be secured stably.
If air cooling is applied instead of accelerated cooling, the cooling rate is too low and the ferrite becomes coarser, and at the same time the bainite fraction decreases, and transformation transformation cannot be sufficiently obtained.
When the accelerated cooling is stopped at a temperature higher than 500 ° C., the accelerated cooling is terminated in the middle of the transformation in the thick plate thickness, so that the transformation strengthening inside the thickness cannot be sufficiently obtained.
In accelerated cooling, it is preferable to secure a water density of 0.3 m 3 / m 2 / min or more to obtain a fine bainite / ferrite structure having both strength and toughness.
The above is a technology that can satisfy the strength on the premise of low Ni under the TMCP condition that places emphasis on the brittle fracture occurrence characteristics, and by this, the needs of (1), (2), and (4) can be satisfied at the same time. .

また、加速冷却後に３５０〜７００℃で５〜６０分の焼戻し熱処理を行ってもよい。これにより、製造コストは上昇するものの、強度や伸び、シャルピー衝撃特性を、高精度で所定の範囲に制御できる。
焼戻し熱処理の温度が３５０℃未満の場合、または焼戻し熱処理の時間が５分未満の場合、焼戻し効果が発揮されない。また、焼戻し熱処理の温度が７００℃超の場合、または焼戻し熱処理の時間が６０分超の場合、焼戻し現象が適正範囲を超えて過剰に発現され、強度低下とシャルピー衝撃特性劣化が著しくなって、適正な機械的性質が得られない。Moreover, you may perform the tempering heat processing for 5 to 60 minutes at 350-700 degreeC after accelerated cooling. Thereby, although the manufacturing cost increases, the strength, elongation, and Charpy impact characteristics can be controlled within a predetermined range with high accuracy.
When the temperature of the tempering heat treatment is less than 350 ° C., or when the time of the tempering heat treatment is less than 5 minutes, the tempering effect is not exhibited. Also, when the temperature of the tempering heat treatment is over 700 ° C., or when the time of the tempering heat treatment is over 60 minutes, the tempering phenomenon is excessively expressed beyond the proper range, and the strength reduction and the Charpy impact property deterioration become remarkable, Appropriate mechanical properties cannot be obtained.

次に、上記（３）のニーズである大入熱溶接ＨＡＺ靭性を満足するための技術について説明する。
本発明の大入熱溶接ＨＡＺ靭性の支配要因は、大別して次の三つである。第一に硬さであり、第二にＭＡ（マルテンサイト・オーステナイト混合相）であり、第三に有効結晶粒径である。
本発明では、硬さとＭＡの両面から、炭素当量Ｃｅｑを０．４２％以下に制限する。炭素当量Ｃｅｑが０．４２％を超えると、ＨＡＺが過剰に硬化すると同時にＭＡが増加し、ＨＡＺが大きく脆化するからである。
さらに、有効Ｂ量（Ｂｅｆ）を０％以下に制御することによって、ＨＡＺにおいてＢ焼入性が発現されることを回避し、硬化とＭＡ増加を抑える。Next, a technique for satisfying the high heat input welding HAZ toughness which is the need of the above (3) will be described.
The controlling factors of the high heat input welding HAZ toughness of the present invention are roughly classified into the following three. First is hardness, second is MA (martensite / austenite mixed phase), and third is effective crystal grain size.
In the present invention, the carbon equivalent Ceq is limited to 0.42% or less in terms of both hardness and MA. This is because if the carbon equivalent Ceq exceeds 0.42%, the HAZ hardens excessively and at the same time the MA increases and the HAZ becomes significantly brittle.
Furthermore, by controlling the effective B amount (Bef) to 0% or less, it is possible to avoid the B hardenability from being expressed in the HAZ, and to suppress hardening and increase in MA.

本発明者等は、硬さの観点から、Ｖ添加の優位性を見出した。また、本発明のようにＨＡＺがベイナイト主体となる場合、Ｖ添加してもＨＡＺは硬化しにくいことを知見した。
つまり、ＣやＭｎなどＶ以外の元素を添加して母材を強化すると、ベイナイト主体のＨＡＺは著しく硬化してＨＡＺは大きく脆化する。これに対して、本発明のようにＶを添加して母材を強化すると、ベイナイト主体のＨＡＺは硬化が抑えられる。この新しい知見に基づくと、Ｖによる母材強度の上昇分を相殺するようにＣやＭｎを低減して低Ｃｅｑ化すれば、ＨＡＺにおいては低Ｃｅｑ化した分だけ硬さが低減するので、ＨＡＺ靭性が向上する。このような、母材とＨＡＺでのＶ硬化挙動の差異を利用したＨＡＺ靭性向上技術は、従来には無かった。The present inventors have found the superiority of V addition from the viewpoint of hardness. Moreover, when HAZ became a bainite main body like this invention, even if it added V, it discovered that HAZ was hard to harden | cure.
That is, when an element other than V, such as C or Mn, is added to strengthen the base material, the bainite-based HAZ is remarkably hardened and the HAZ is greatly embrittled. On the other hand, when the base material is strengthened by adding V as in the present invention, the hardening of HAZ mainly composed of bainite is suppressed. Based on this new knowledge, if C and Mn are reduced to lower the Ceq so as to offset the increase in the base metal strength due to V, the hardness will decrease in HAZ by the reduced Ceq, so HAZ Toughness is improved. Conventionally, there has been no HAZ toughness improving technique using the difference in V-curing behavior between the base material and HAZ.

本発明では、ＭＡの観点から、可能な限りＳｉを低減する必要がある。
また、本発明のＴＭＣＰ条件では、Ｎｂは母材材質への貢献が小さいにも関わらずＭＡ生成を助長する。本発明の比較的高いＣｅｑ範囲では、Ｍｏは高価であるにも関わらずＭＡ生成を助長する。従って、ＮｂとＭｏは本発明においては可能な限り低減する必要がある。In the present invention, it is necessary to reduce Si as much as possible from the viewpoint of MA.
In addition, under the TMCP condition of the present invention, Nb promotes MA generation despite its small contribution to the base material. In the relatively high Ceq range of the present invention, Mo facilitates MA formation despite being expensive. Therefore, Nb and Mo must be reduced as much as possible in the present invention.

本発明では、有効結晶粒径の観点から、二つのＨＡＺ組織微細化技術を適用する。
第一は、γ中のＢ析出物とＶ析出物を変態核として同時に利用することである。上記式（２）で表される有効Ｂ量｛Ｂｅｆ（％）｝が０％以下となるようにＮ量を適性に高めることによって、大入熱溶接の冷却中にγ粒界やγ粒内にＢＮとＶＮ、Ｖ（Ｃ，Ｎ）が析出し、これらの単独あるいは複合の粒子がフェライトのみならずベイナイトの変態核としても有効に作用し、ＨＡＺ組織を微細化する。
さらにＨＡＺ組織を微細化する第二の技術は、ＣａやＭｇの適正添加によって微細な酸化物や硫化物を多数分散させ、γ粒成長をピン止め効果によって抑制することで、ベイナイトのパケットを微細化する。微細な酸化物や硫化物の一部には、Ｂ析出物やＶ析出物が複合析出し、ピン止め粒子に変態核機能が付加されることで、γ粒界から変態するベイナイトをより一層微細化する効果もある。
以上のようなＨＡＺ組織微細化技術は、結果的にＨＡＺの焼入性を低めるので、硬さとＭＡを低減する観点からも貢献する。第一の技術によって−２０℃のシャルピー吸収エネルギーを確保し、これに第二の技術を組み合わせることでＨＡＺ組織を極限まで微細化すれば、−４０℃のシャルピー吸収エネルギーを確保できる可能性がある。
以上説明した硬さ低減、ＭＡ低減、及びＨＡＺ組織微細化の施策を通じて、本発明の大入熱溶接ＨＡＺは高いｖＥ（−２０℃）を達成することができる。これにより、上記（１）、（２）、及び（４）に加えて、（３）のニーズを満たすことが可能となる。In the present invention, two HAZ structure refinement techniques are applied from the viewpoint of effective crystal grain size.
The first is to simultaneously use B precipitates and V precipitates in γ as transformation nuclei. By appropriately increasing the N amount so that the effective B amount {Bef (%)} represented by the above formula (2) is 0% or less, the γ grain boundary and the γ grain inside the high heat input welding are cooled. BN, VN, and V (C, N) are precipitated on these particles, and these single or composite particles effectively act not only as ferrite but also as a transformation nucleus of bainite, thereby refining the HAZ structure.
In addition, the second technology for refining the HAZ structure is to finely disperse bainite packets by dispersing a large number of fine oxides and sulfides by appropriate addition of Ca and Mg, and suppressing γ grain growth by the pinning effect. Turn into. Part of fine oxides and sulfides are combined with B precipitates and V precipitates, and the transformation nucleus function is added to the pinning particles, so that the bainite transformed from the γ grain boundaries is made even finer. There is also an effect that.
Since the HAZ structure refinement technique as described above lowers the hardenability of HAZ as a result, it contributes from the viewpoint of reducing hardness and MA. If Charpy absorption energy of −20 ° C. is secured by the first technique and the HAZ structure is refined to the limit by combining this with the second technique, Charpy absorption energy of −40 ° C. may be secured. .
Through the measures for reducing the hardness, reducing the MA, and refining the HAZ structure as described above, the high heat input welding HAZ of the present invention can achieve a high vE (−20 ° C.). Thereby, in addition to the above (1), (2), and (4), it is possible to satisfy the needs of (3).

＜化学成分組成（厚手高強度鋼板）＞
本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板は、上述したような、（１）大きな板厚での高い強度、（２）良好な脆性破壊伝播停止特性、（３）良好な大入熱溶接ＨＡＺ靭性、及び（４）低い製造コスト等のニーズを満足するため、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含み、強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量を、下記式（５）で表される量としたとき、下記式（６）で表される、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、さらに、下記式（７）で表される炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たし、板厚が５０〜８０ｍｍであり、降伏強度が３９０〜４６０ＭＰａ級で、引張強度が５１０〜５７０ＭＰａ級であり、脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００が−１０℃以下であり、溶接入熱量が２０ｋＪ／ｍｍ以上の大入熱溶接部のＨＡＺ靭性の指標であるシャルピー衝撃吸収エネルギーｖＥ（−２０℃）が４７Ｊ以上とされている。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（５）｛但し、式（５）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（６）｛但し、式（６）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（７）
なお、上記式（５）〜（７）の各式において、式（５）は上記式（１）と共通の式であり、また、式（６）は上記式（２）、式（７）は上記式（３）とそれぞれ共通の式である。<Chemical component composition (thick high-strength steel plate)>
The thick high-strength steel plate excellent in brittle fracture propagation stop characteristics and high heat input welding heat-affected zone toughness according to the present invention is as described above, (1) high strength at large plate thickness, and (2) good brittle fracture. In order to satisfy the needs such as propagation stop characteristics, (3) good high heat input welding HAZ toughness, and (4) low production cost, C: 0.05 to 0.12%, Si: 0.00. 3% or less, Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al : 0.001 to 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, and O: 0.004% or less, the balance being iron and inevitable The amount of residual oxygen that contains impurities and remains after deoxidation with a strong deoxidation element and can be deoxidized with Ti, which is a weak deoxidation element, is When the amount represented by (5) is used, the calculated value of the B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation represented by the following formula (6) is 0% or less. Furthermore, the carbon equivalent Ceq represented by the following formula (7) satisfies the range of 0.32 to 0.42%, the plate thickness is 50 to 80 mm, the yield strength is 390 to 460 MPa class, the tensile strength is High heat input welding with strength of 510-570 MPa class, temperature T kca = 6000 at which brittle fracture propagation stop characteristic Kca is 6000 N / mm 1.5 is −10 ° C. or less, and welding heat input is 20 kJ / mm or more. Charpy impact absorption energy vE (−20 ° C.) that is an index of HAZ toughness of the part is 47 J or more.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (5) {However, in the formula (5), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (6) {However, in the formula (6), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (7)
In each of the above formulas (5) to (7), formula (5) is a common formula with formula (1), and formula (6) is formula (2) and formula (7). Are common to the above equation (3).

また、本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板は、上記化学成分組成において、上記各元素のうち、Ｓの含有量の下限を０．０００５％、Ｏの含有量の下限を０．００１％とすることができる。さらに、必要に応じて、Ｃａ：０．０００３〜０．００４％、Ｍｇ：０．０００３〜０．００４％、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、Ｎｂ：０．００３〜０．０３％、ＲＥＭ：０．０００３〜０．０２％、及びＺｒ：０．０００３〜０．０２％のうちの１種または２種以上を選択的に含有する構成とすることができる。
以下に、本発明における鋼（厚手高強度鋼板）の化学成分についての限定理由を説明する。Further, the thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention has a lower limit of the content of S of the above-mentioned elements in the above chemical composition composition. The lower limit of 0005% and O content can be 0.001%. Furthermore, if necessary, Ca: 0.0003 to 0.004%, Mg: 0.0003 to 0.004%, Ni: 0.01 to 1%, Cu: 0.01 to 1%, Cr: 0 0.01-1%, Mo: 0.01-0.5%, Nb: 0.003-0.03%, REM: 0.0003-0.02%, and Zr: 0.0003-0.02% It can be set as the structure which contains selectively 1 type (s) or 2 or more types.
Below, the reason for limitation about the chemical component of the steel (thick high-strength steel plate) in this invention is demonstrated.

「Ｃ：炭素」０．０５〜０．１２％
Ｃは、強度向上のために重要な元素である。低温加熱、低温圧延を徹底したＴＭＣＰ型厚手鋼板において、所定の強度を安定確保するために、０．０５％以上のＣを添加する必要がある。また、後述する理由から、本発明ではＮｂ、Ｎｉ、及びＭｏの添加量を必要最小限に抑える必要があるので、これらの元素を増加して高強度化することは困難である。従って、Ｃは非常に重要な強化元素である。さらに、Ｃは大入熱ＨＡＺにおけるＶ（Ｃ，Ｎ）変態核の析出を促す効果もある。しかしながら、良好なＨＡＺ靭性を安定確保するためには、Ｃを０．１２％以下に抑えることが必要であり、ＨＡＺ靭性を高めるためには０．１０％以下とすることが好ましい。“C: Carbon” 0.05 to 0.12%
C is an important element for improving the strength. In a TMCP type thick steel plate thoroughly subjected to low temperature heating and low temperature rolling, it is necessary to add 0.05% or more of C in order to ensure a predetermined strength stably. In addition, for the reason described later, in the present invention, it is necessary to suppress the addition amount of Nb, Ni, and Mo to a necessary minimum, so it is difficult to increase the strength by increasing these elements. Therefore, C is a very important strengthening element. Furthermore, C also has an effect of promoting the precipitation of V (C, N) transformation nuclei in the high heat input HAZ. However, in order to stably secure good HAZ toughness, it is necessary to suppress C to 0.12% or less, and in order to increase HAZ toughness, it is preferable to set it to 0.10% or less.

「Ｓｉ：ケイ素」０．３％以下
Ｓｉは、脱酸作用を有するが、強力な脱酸元素であるＡｌが十分に添加されている場合には不要である。Ｓｉは、母材を強化する作用もあるが、他の元素に比べるとその効果は相対的に小さい。また、比較的高い炭素当量Ｃｅｑが必要となる本発明の大入熱溶接ＨＡＺでは、ＳｉはＭＡ生成を助長する危険性が高いため、０．３％以下に抑える必要があり、ＨＡＺ靭性の観点からＳｉの添加量を極力低くなるように、０．２０％以下とすることが好ましい。強度確保と脱酸をおこなうために、Ｓｉを０．０１％以上添加することが望ましい。“Si: silicon” 0.3% or less Si has a deoxidizing action, but is unnecessary when Al, which is a strong deoxidizing element, is sufficiently added. Si also has the effect of strengthening the base material, but its effect is relatively small compared to other elements. In addition, in the high heat input welding HAZ of the present invention that requires a relatively high carbon equivalent Ceq, since Si has a high risk of promoting the formation of MA, it is necessary to suppress it to 0.3% or less. From the viewpoint of HAZ toughness From the above, it is preferable to make the addition amount of Si 0.20% or less so that the addition amount of Si is as low as possible. In order to ensure strength and perform deoxidation, it is desirable to add 0.01% or more of Si.

「Ｍｎ：マンガン」１〜２％
Ｍｎは、経済的に強度を確保するために１％以上の添加量が必要であり、１．４０％以上とすることが好ましい。但し、２％を超えてＭｎを添加すると、スラブの中心偏析の有害性が顕著となるうえ、大入熱溶接ＨＡＺの硬化とＭＡ生成を助長して脆化させるため、これを上限とする。この脆化を防止するために、Ｍｎを１．６０％以下に制限することが好ましい。"Mn: Manganese" 1-2%
Mn requires an addition amount of 1% or more in order to ensure strength economically, and is preferably 1.40% or more. However, if Mn is added in excess of 2%, the hazard of center segregation of the slab becomes remarkable, and the hardening and MA formation of the high heat input weld HAZ are promoted and become brittle, so this is made the upper limit. In order to prevent this embrittlement, it is preferable to limit Mn to 1.60% or less.

「Ｐ：リン」０．０１５％以下
Ｐは、不純物元素であり、良好な脆性破壊伝播停止特性と大入熱溶接ＨＡＺ靭性を安定的に確保するために、０．０１５％以下に低減する必要がある。ＨＡＺ靭性を高めるためには、０．０１０％以下とすることが好ましい。“P: Phosphorus” 0.015% or less P is an impurity element, and it is necessary to reduce it to 0.015% or less in order to stably secure good brittle fracture propagation stopping characteristics and high heat input HAZ toughness. There is. In order to increase the HAZ toughness, the content is preferably 0.010% or less.

「Ｓ：硫黄」０．０００５〜０．００５％
Ｓは、０．００５％以下に抑える必要がある。Ｓが０．００５％を超えると、硫化物の一部が粗大化して破壊起点として有害性をもたらし、母材と大入熱溶接ＨＡＺの靭性が劣化する。この有害性をより少なくするために、０．００３％以下とすることが好ましい。一方で、ＨＡＺのピン止め効果を利用する際には、Ｓは０．０００５％以上確保する必要がある。その理由は、ＨＡＺの溶融線近傍において、ＨＡＺ靭性を高めるためにＣａやＭｇの適正添加によって微細な硫化物を多数分散させ、ピン止め効果を強化してγ細粒化を図るためである。Ｓが０．０００５％未満の場合、硫化物個数が不足して十分なピン止め効果が得られない。“S: sulfur” 0.0005 to 0.005%
S needs to be suppressed to 0.005% or less. When S exceeds 0.005%, a part of the sulfide is coarsened to cause harmfulness as a fracture starting point, and the toughness of the base metal and the high heat input welding HAZ deteriorates. In order to reduce this harmfulness, it is preferable to make it 0.003% or less. On the other hand, when utilizing the pinning effect of HAZ, S needs to be secured at 0.0005% or more. The reason is that, in the vicinity of the HAZ melting line, a large amount of fine sulfides are dispersed by appropriate addition of Ca and Mg in order to increase the HAZ toughness, and the pinning effect is strengthened to achieve γ fine graining. When S is less than 0.0005%, the number of sulfides is insufficient and a sufficient pinning effect cannot be obtained.

「Ｂ：ボロン（ホウ素）」０．０００３〜０．００３％
Ｂは、本発明の特徴的な元素である。既に詳述したように、本発明では母材と大入熱溶接ＨＡＺの両方において、γ中に固溶Ｂを存在させずに全てのＢをＢＮとして析出させ、Ｂ焼入性が発現しないように、上記式（２）で表される有効Ｂ量（Ｂｅｆ）の算出値を０％以下に制御する。γ中に析出させたＢＮは変態核として作用し、ＨＡＺの組織微細化、硬さ低減、及びＭＡ低減を通じて靭性を高める。そのためには、Ｂを０．０００３％以上添加する必要がある。但し、０．００３％を超えてＢを添加すると、粗大なＢ析出物が生成してＨＡＺ靭性が劣化するため、これを上限とする。ＨＡＺ靭性を安定的に確保するためには、０．００２０％以下とすることが好ましい。“B: Boron” 0.0003 to 0.003%
B is a characteristic element of the present invention. As already described in detail, in the present invention, in both the base material and the high heat input welding HAZ, all B is precipitated as BN without the presence of solid solution B in γ so that the B hardenability does not appear. In addition, the calculated value of the effective B amount (Bef) represented by the above formula (2) is controlled to 0% or less. BN precipitated in γ acts as a transformation nucleus and enhances toughness through HAZ microstructure refinement, hardness reduction, and MA reduction. For that purpose, it is necessary to add 0.0003% or more of B. However, if B is added in excess of 0.003%, coarse B precipitates are generated and the HAZ toughness deteriorates, so this is the upper limit. In order to stably secure the HAZ toughness, the content is preferably 0.0020% or less.

「Ｖ：バナジウム」０．０１〜０．１５％
Ｖは、本発明の特徴的な元素である。既に詳述したように、Ｖは本発明のＴＭＣＰ条件において母材を効果的に強化する。その一方で、Ｖは本発明の大入熱溶接ＨＡＺにおいて硬化やＭＡ増加を抑えると同時に、γ中に析出させたＶＮやＶ（Ｃ，Ｎ）は変態核として作用し、ＨＡＺ組織を微細化して靭性を高める。この効果を発揮するためには、０．０１％以上のＶが必要である。しかしながら、Ｖが０．１５％を超えると、ＨＡＺの組織微細化効果が飽和すると同時にＨＡＺの硬化が著しくなるので、ＨＡＺ靭性が劣化する。従って、０．１５％がＶの上限であり、０．１０％以下とすることが好ましい。“V: Vanadium” 0.01 to 0.15%
V is a characteristic element of the present invention. As already described in detail, V effectively strengthens the base material under the TMCP conditions of the present invention. On the other hand, V suppresses hardening and MA increase in the high heat input welding HAZ of the present invention, and at the same time, VN and V (C, N) precipitated in γ act as transformation nuclei to refine the HAZ structure. Increase toughness. In order to exert this effect, 0.01% or more of V is necessary. However, if V exceeds 0.15%, the effect of refining the HAZ structure is saturated, and at the same time, the hardening of the HAZ becomes remarkable, so that the HAZ toughness deteriorates. Therefore, 0.15% is the upper limit of V, and is preferably 0.10% or less.

「Ａｌ：アルミニウム」０．００１〜０．１％
Ａｌは、脱酸を担い、Ｏを低減して鋼の清浄度を高めるために必要である。Ａｌ以外のＳｉ、Ｔｉ、Ｃａ、Ｍｇ、ＲＥＭ、Ｚｒ等も脱酸作用があるが、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（酸素）を０．００４％以下に抑えることは難しい。但し、Ａｌが０．１％を超えるとアルミナ系粗大酸化物がクラスター化する傾向を強め、製鋼ノズルつまりが発生したり、破壊起点としての有害性が顕在化するため、これを上限とする。有害性をより少なくするためには、Ａｌを０．０６０％以下とすることが好ましい。“Al: Aluminum” 0.001 to 0.1%
Al is necessary for carrying out deoxidation, reducing O, and increasing the cleanliness of steel. Si, Ti, Ca, Mg, REM, Zr, etc. other than Al also have a deoxidizing action, but even when these elements are added, O (oxygen) is stable without 0.001% or more of Al. ) To 0.004% or less is difficult. However, if Al exceeds 0.1%, the tendency of the alumina-based coarse oxide to be clustered is strengthened, and a steelmaking nozzle is clogged, or the harmfulness as a fracture starting point becomes obvious, so this is the upper limit. In order to further reduce the harmfulness, Al is preferably made 0.060% or less.

「Ｔｉ：チタン」０．００５〜０．０２％、「Ｎ：窒素」０．００２〜０．０１％、及び「有効Ｂ量：Ｂｅｆ（％）」０％以下（上記式（２）の算出値）
Ｔｉは、Ｎと結合してＴｉＮを形成し、スラブ再加熱時と大入熱溶接ＨＡＺでピン止め効果に貢献し、γ細粒化に寄与する結果、母材やＨＡＺの組織を微細化して靭性を高める。そして、ＴｉＮを形成した残りのＮはＢと結合してＢＮを形成し、γ中に固溶Ｂを存在させずに全てのＢをＢＮとして析出させ、Ｂ焼入性が発現しないようにする。
以上のような効果を同時に発揮するためには、Ｔｉを０．００５〜０．０２％、Ｎを０．００２〜０．０１％、及び上記式（２）で表される有効Ｂ量（Ｂｅｆ）の算出値を０％以下とする必要がある。
ＴｉとＮが、それぞれ０．００５％、０．００２％に満たないと、ＴｉＮによるピン止め効果が十分に発揮されず、母材とＨＡＺの靭性が劣化する。ＴｉとＮがそれぞれ０．０２％、０．０１％を超えると、ＴｉＣ析出や固溶Ｎ増加によって母材とＨＡＺの靭性が劣化する。よりＨＡＺ靭性を高くするために、ＴｉとＮをそれぞれ０．０１５％、０．００７％以下とすることが望ましい。さらに、ＴｉとＮが適正範囲にあっても、有効Ｂ量が０％を超えると、γ中の固溶Ｂの量が増加してＢ焼入性が発現し、母材強度のばらつきやＨＡＺの硬化（脆化）をもたらす。“Ti: titanium” 0.005 to 0.02%, “N: nitrogen” 0.002 to 0.01%, and “effective B amount: Bef (%)” 0% or less (calculation of the above formula (2) value)
Ti combines with N to form TiN, contributes to pinning effect during slab reheating and high heat input welding HAZ, and contributes to γ grain refinement, resulting in refined base metal and HAZ structure Increase toughness. And the remaining N which formed TiN couple | bonds with B, forms BN, makes all B precipitate as BN, without making solid solution B exist in (gamma), and B hardenability is not expressed. .
In order to exhibit the above effects at the same time, 0.005 to 0.02% of Ti, 0.002 to 0.01% of N, and an effective B amount represented by the above formula (2) (Bef ) Must be 0% or less.
If Ti and N are less than 0.005% and 0.002%, respectively, the pinning effect by TiN is not sufficiently exhibited, and the toughness of the base material and the HAZ deteriorates. When Ti and N exceed 0.02% and 0.01%, respectively, the toughness of the base material and the HAZ deteriorates due to TiC precipitation and solute N increase. In order to further increase the HAZ toughness, it is desirable that Ti and N be 0.015% and 0.007% or less, respectively. Furthermore, even if Ti and N are in an appropriate range, if the amount of effective B exceeds 0%, the amount of solute B in γ increases and B hardenability is developed, resulting in variations in base material strength and HAZ. Causes hardening (embrittlement).

以下に、有効Ｂ量（Ｂｅｆ）の考え方を説明する。
化学成分として添加されたＴｉは、溶鋼中の脱酸で消費される場合があり（低Ａｌの場合に起こりやすい）、脱酸後に残ったＴｉが凝固後のγ中でＴｉＮを形成する。この際、Ｔｉに対してＮが過剰であると、ＴｉＮを形成した後に残ったＮがＢの一部と結合してＢＮを形成する。そして、ＢＮを形成した残りのＢが固溶Ｂとして焼入性を発現してしまう。この焼入性に寄与するγ中の固溶Ｂ量を本発明では有効Ｂ量Ｂｅｆ（％）として扱う。The concept of the effective B amount (Bef) will be described below.
Ti added as a chemical component may be consumed by deoxidation in molten steel (prone to occur in the case of low Al), and Ti remaining after deoxidation forms TiN in γ after solidification. At this time, if N is excessive with respect to Ti, N remaining after forming TiN is combined with a part of B to form BN. And the remaining B which formed BN will express hardenability as the solid solution B. In the present invention, the amount of dissolved B in γ that contributes to the hardenability is treated as an effective B amount Bef (%).

各元素の添加量、熱力学的な反応順序、及び生成物質の化学量論組成に基づいた有効Ｂ量Ｂｅｆの計算方法について以下に説明する。
まず、脱酸力の高い順に、Ｃａ、Ｍｇ、ＲＥＭ（希土類元素）、Ｚｒ、及びＡｌがＯと結合すると仮定する。この際の脱酸生成物として、ＣａＯ、ＭｇＯ、ＲＥＭ２Ｏ３、ＺｒＯ２、及びＡｌ２Ｏ３を仮定して、脱酸されるＯ量を計算する。
Ｔｉよりも脱酸力の強いこれらの元素によって脱酸が完了しない場合、これらの強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉによって脱酸され得る残存酸素量ＯＴｉ（％）を、下記式（１）で表される量とした時、次式｛ＯＴｉ（％）＞０｝を満たす。
ＯＴｉ（％） ＝ Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（１）
但し、上記式（１）において、不可避的不純物扱いの成分元素も計算に含める。A method for calculating the effective B amount Bef based on the addition amount of each element, the thermodynamic reaction sequence, and the stoichiometric composition of the product will be described below.
First, it is assumed that Ca, Mg, REM (rare earth element), Zr, and Al are combined with O in descending order of deoxidizing power. Assuming CaO, MgO, REM 2 O 3 , ZrO 2 , and Al 2 O 3 as deoxidation products at this time, the amount of O to be deoxidized is calculated.
In the case where deoxidation is not completed by these elements having a stronger deoxidizing power than Ti, the residual oxygen amount O Ti (%) that remains after deoxidation by these strong deoxidizing elements and can be deoxidized by Ti, which is a weak deoxidizing element ) Is the amount represented by the following formula (1), the following formula {O Ti (%)> 0} is satisfied.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1)
However, in the above formula (1), component elements that are treated as inevitable impurities are also included in the calculation.

この場合、残ったＯ（＝ＯＴｉ）をＴｉが脱酸することになる。そこで、Ｔｉ２Ｏ３を仮定して、脱酸で消費されるＴｉを差し引いた残りのＴｉは、Ｔｉ−２ＯＴｉ≧０．００５（％）で表され、この値が０．００５％以上になる必要がある。ここで、脱酸で消費されるＴｉを差し引いた残りのＴｉが０．００５％以上必要なのは、上述したように、本発明に必要なＴｉＮを確保するためである。
脱酸で消費されるＴｉを差し引いた残りのＴｉが０．００５％未満であると、ＴｉＮによるピン止め効果が十分に発揮されず、厚手母材と大入熱溶接ＨＡＺ靭性が劣化する。In this case, Ti deoxidizes the remaining O (= O Ti ). Therefore, assuming Ti 2 O 3 , the remaining Ti after subtracting Ti consumed in deoxidation is expressed as Ti-2O Ti ≧ 0.005 (%), and this value is 0.005% or more. Need to be. Here, the reason why 0.005% or more of the remaining Ti after subtracting Ti consumed in deoxidation is necessary is to secure TiN necessary for the present invention as described above.
If the remaining Ti after subtracting Ti consumed by deoxidation is less than 0.005%, the pinning effect by TiN is not sufficiently exhibited, and the thick base material and the high heat input weld HAZ toughness deteriorate.

また、脱酸で残った０．００５％以上のＴｉがＴｉＮを形成し、Ｎが残る場合は下記式が正の値になり、Ｎが残らない場合は下記式が０又は負の値になる。
Ｎ−０．２９（Ｔｉ−２ＯＴｉ） ＞ ０ ：Ｎが残る場合
Ｎ−０．２９（Ｔｉ−２ＯＴｉ） ≦ ０ ：Ｎが残らない場合Further, 0.005% or more of Ti remaining after deoxidation forms TiN, and when N remains, the following formula becomes a positive value, and when N does not remain, the following formula becomes 0 or a negative value. .
N-0.29 (Ti-2O Ti )> 0: If N remains N-0.29 (Ti-2O Ti ) ≦ 0: If N does not remain

また、上記式｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝が正の値となってＮが残る場合は、Ｂの一部がＢＮとして消費されるので、下記式（２）によって有効Ｂ量Ｂｅｆが算出される。
Ｂｅｆ（％） ＝ Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）
但し、上記式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、式｛Ｔｉ−２ＯＴｉ≧０．００５（％）｝を満たすものとする。さらに、式｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）｝のときは、式｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０｝とする。Further, when the above expression {N-0.29 (Ti-2O Ti )} becomes a positive value and N remains, a part of B is consumed as BN. The quantity Bef is calculated.
Bef (%) = B-0.77 {N-0.29 (Ti-2O Ti)} ··· (2)
However, in the above formula (2), when O Ti ≦ 0, O Ti = 0. Further, when O Ti > 0, the expression {Ti-2O Ti ≧ 0.005 (%)} is satisfied. Further, the formula {N-0.29 (Ti-2O Ti) ≦ 0 ( provided that when the O Ti ≦ 0, O Ti = 0)} When the equation {N-0.29 (Ti-2O Ti) = 0}.

また、式｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝が０または負の値となってＮが残らない場合は、有効Ｂ量Ｂｅｆは、Ｂｅｆ（％）＝Ｂで表される量となる。Further, when the expression {N−0.29 (Ti−2O Ti )} is 0 or a negative value and N does not remain, the effective B amount Bef is the amount represented by Bef (%) = B. Become.

次に、上述した残存酸素量ＯＴｉの式におけるＣａ、Ｍｇ、ＲＥＭ、Ｚｒ、及びＡｌの係数について述べると、溶鋼中での脱酸反応（酸化反応）による生成物（酸化物）としてＣａＯ、ＭｇＯ、ＲＥＭ２Ｏ３、ＺｒＯ２、Ａｌ２Ｏ３を仮定し、これらの酸化物として存在するＯ量を質量％で計算する。例えば、ＣａＯの場合、原子量はＣａが４０でＯが１６であるから、Ｃａの質量％に対して１６／４０＝０．４のＯが結合する。Ａｌ２Ｏ３であれば、原子量はＡｌが２７でＯが１６であるから、Ａｌの質量％に対して（１６×３）／（２７×２）＝０．８９のＯが結合する。以下同様の計算概念として、上述のＯＴｉ式の各元素の係数（０．６６：Ｍｇ、０．１７：ＲＥＭ、及び０．３５：Ｚｒ）を規定した。Next, the coefficients of Ca, Mg, REM, Zr, and Al in the residual oxygen amount O Ti formula described above will be described. As a product (oxide) by deoxidation reaction (oxidation reaction) in molten steel, CaO, Assuming MgO, REM 2 O 3 , ZrO 2 , Al 2 O 3 , the amount of O present as these oxides is calculated in mass%. For example, in the case of CaO, since the atomic weight is 40 for Ca and 16 for O, O of 16/40 = 0.4 is bonded to the mass% of Ca. In the case of Al 2 O 3 , since the atomic weight is 27 for Al and 16 for O, O of (16 × 3) / (27 × 2) = 0.89 is bonded to the mass% of Al. Hereinafter, as the same calculation concept, the coefficients (0.66: Mg, 0.17: REM, and 0.35: Zr) of each element of the above O Ti formula were defined.

また、有効Ｂ量Ｂｅｆの導出式概念を、低温側から高温側に遡って示すと以下のようになる。
有効Ｂ量Ｂｅｆ（％） ＝ 成分Ｂ量−Ｂ ａｓ ＢＮ
→ Ｂ ａｓ ＢＮ ＝ ０．７７（Ｎ − Ｎ ａｓ ＴｉＮ）
→ Ｎ ａｓ ＴｉＮ ＝ ０．２９（Ｔｉ − Ｔｉ ａｓ Ｔｉ２Ｏ３）
→ Ｔｉ ａｓ Ｔｉ２Ｏ３ ＝ ２（Ｏ − Ｏ ａｓ ＣａＯ − Ｏ ａｓ ＭｇＯ − Ｏ ａｓ ＲＥＭ２Ｏ３ − Ｏ ａｓ ＺｒＯ２ − Ｏ ａｓ Ａｌ２Ｏ３）
→ Ｏ ａｓ ＣａＯ ＝ ０．４Ｃａ
→ Ｏ ａｓ ＭｇＯ ＝ ０．６６Ｍｇ
→ Ｏ ａｓ ＲＥＭ２Ｏ３ ＝ ０．１７ＲＥＭ
→ Ｏ ａｓ ＺｒＯ２ ＝ ０．３５Ｚｒ
→ Ｏ ａｓ Ａｌ２Ｏ３ ＝ ０．８９ＡｌFurther, the concept of the derivation formula for the effective B amount Bef is shown as follows from the low temperature side to the high temperature side.
Effective B amount Bef (%) = component B amount−B as BN
→ B as BN = 0.77 (N-N as TiN)
→ N as TiN = 0.29 (Ti - Ti as Ti 2 O 3)
→ Ti as Ti 2 O 3 = 2 (O - O as CaO - O as MgO - O as REM 2 O 3 - O as ZrO 2 - O as Al 2 O 3)
→ O as CaO = 0.4Ca
→ O as MgO = 0.66Mg
→ O as REM 2 O 3 = 0.17 REM
→ O as ZrO 2 = 0.35Zr
→ O as Al 2 O 3 = 0.89Al

次に、有効Ｂ量Ｂｅｆの導出式概念を、高温側から低温側への反応順に示すと以下のようになる。すなわち、製鋼での精錬→凝固工程において、以下の順で反応する。   Next, the derivation concept of the effective B amount Bef is shown as follows in the order of reaction from the high temperature side to the low temperature side. That is, in the refining-> solidification process in steelmaking, the reaction occurs in the following order.

［１］液相（溶鋼中）での脱酸反応（１６００℃付近）
Ｏとの化学的親和力の強い順にＣａＯ→ＭｇＯ→ＲＥＭ２Ｏ３→ＺｒＯ２→Ａｌ２Ｏ３の反応が生じ、溶鋼中の溶存Ｏが減少していく。これで脱酸が完了する場合は、ＯＴｉ≦０で表される。脱酸が完了せずに溶存Ｏが残る場合は、ＯＴｉ＞０、Ｔｉ−２ＯＴｉ≧０．００５（％）で表され、Ａｌより弱脱酸元素であるＴｉがＴｉ２Ｏ３として脱酸に寄与し、成分Ｔｉから脱酸で消費されたＴｉ ａｓ Ｔｉ２Ｏ３を差し引いた残りのＴｉが０．００５％以上となる。[1] Deoxidation reaction in liquid phase (in molten steel) (around 1600 ° C)
The reaction of CaO → MgO → REM 2 O 3 → ZrO 2 → Al 2 O 3 occurs in the order of strong chemical affinity with O, and the dissolved O in the molten steel decreases. When deoxidation is completed by this, it is represented by O Ti ≦ 0. When dissolved O remains without completing deoxidation, it is represented by O Ti > 0, Ti-2O Ti ≧ 0.005 (%), and Ti, which is a weaker deoxidizing element than Al, is desorbed as Ti 2 O 3. The remaining Ti that contributes to the acid and subtracts the Ti as Ti 2 O 3 consumed by deoxidation from the component Ti becomes 0.005% or more.

［２］固相（凝固γ中）での脱窒反応（１３００℃付近〜８００℃付近）
Ｎとの化学的親和力の強い順にＴｉＮ→ＢＮ→ＡｌＮの反応が生じ、固相γ中の固溶Ｎが減少していく。まず、脱酸で消費された残りのＴｉが脱窒反応を起こす。これで脱窒が完了する場合は、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０で表され、γ中に固溶Ｎが存在しないので、ＢはＢＮを形成せずに全てが固溶Ｂとして存在する。一方、Ｔｉによって脱窒が完了せず、固溶Ｎが残る場合は、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＞０で表され、Ｂの一部がＢＮを生成して残りが固溶Ｂとなる。[2] Denitrification reaction in solid phase (in solidification γ) (around 1300 ° C to 800 ° C)
The reaction of TiN → BN → AlN occurs in the order of strong chemical affinity with N, and the solid solution N in the solid phase γ decreases. First, the remaining Ti consumed by deoxidation causes a denitrification reaction. When denitrification is completed in this way, N−0.29 (Ti−2O Ti ) ≦ 0, and since solute N does not exist in γ, B does not form BN but is completely dissolved. Exists as B. On the other hand, when denitrification is not completed by Ti and solid solution N remains, it is represented by N-0.29 (Ti-2O Ti )> 0, and a part of B generates BN and the remaining is solid solution. B.

一方、Ｔｉよりも脱酸力の強い元素によって脱酸が完了する場合には、下記式を満たす。
ＯＴｉ≦０
この場合、Ｔｉは脱酸では消費されない。ＴｉがＴｉＮを形成し、Ｎが残る場合は下記式を満たす。
Ｎ−０．２９Ｔｉ＞０
この際の有効Ｂ量Ｂｅｆは下記式で計算される。
Ｂｅｆ（％）＝Ｂ−０．７７（Ｎ−０．２９Ｔｉ）
ＴｉがＴｉＮを形成し、Ｎが残らない場合は下記式を満たす。
Ｎ−０．２９Ｔｉ≦０
この際の有効Ｂ量Ｂｅｆは下記式で計算される。
Ｂｅｆ（％）＝ＢOn the other hand, when deoxidation is completed by an element having a stronger deoxidizing power than Ti, the following formula is satisfied.
O Ti ≦ 0
In this case, Ti is not consumed by deoxidation. When Ti forms TiN and N remains, the following formula is satisfied.
N-0.29Ti> 0
The effective B amount Bef at this time is calculated by the following equation.
Bef (%) = B−0.77 (N−0.29Ti)
When Ti forms TiN and N does not remain, the following formula is satisfied.
N−0.29Ti ≦ 0
The effective B amount Bef at this time is calculated by the following equation.
Bef (%) = B

なお、上記各式において、式（Ｎ−０．２９Ｔｉ）における０．２９Ｔｉは、Ｎ ａｓ ＴｉＮを意味する。ここで、原子量はＴｉが４８でＮが１４であるから、Ｔｉ（正確には脱酸で消費されたＴｉを差し引いた残りのＴｉ）の質量％に対して１４／４８＝０．２９のＮが結合する。また、Ｎ−０．２９Ｔｉ≦０であれば、Ｎは全てＴｉＮで固定され、γ素地中に固溶Ｎは存在しない。一方、Ｎ−０．２９Ｔｉ＞０ならば、γ素地中にはＴｉＮの他に固溶Ｎが存在するので、この固溶Ｎは、Ｂと結合してＢＮを生成し、有効Ｂ量を減少させる。   In the above formulas, 0.29Ti in the formula (N-0.29Ti) means N as TiN. Here, since the atomic weight is 48 for Ti and 14 for N, N of 14/48 = 0.29 with respect to the mass% of Ti (exactly Ti after subtracting Ti consumed in deoxidation). Join. If N−0.29Ti ≦ 0, all N is fixed with TiN, and no solute N exists in the γ substrate. On the other hand, if N−0.29Ti> 0, solid solution N exists in addition to TiN in the γ substrate, so this solid solution N combines with B to produce BN, reducing the effective B amount. Let

「Ｏ：酸素」０．００１〜０．００４％以下
Ｏは、０．００４％以下に抑える必要がある。Ｏが０．００４％を超えると、酸化物の一部が粗大化して破壊起点として有害性をもたらし、母材と大入熱溶接ＨＡＺの靭性が劣化する。一方で、ＨＡＺのピン止め効果を利用する際には、Ｏは０．００１％以上確保する必要がある。その理由は、ＨＡＺの溶融線近傍において、ＨＡＺ靭性を高めるためにＣａやＭｇの適正添加によって微細な酸化物を多数分散させ、ピン止め効果を強化してγ細粒化を図るためである。Ｏが０．００１％未満の場合、酸化物個数が不足して十分なピン止め効果が得られなくなる虞がある。“O: Oxygen” 0.001 to 0.004% or less O needs to be suppressed to 0.004% or less. When O exceeds 0.004%, a part of the oxide is coarsened to cause harmfulness as a fracture starting point, and the toughness of the base material and the high heat input welding HAZ is deteriorated. On the other hand, when utilizing the pinning effect of HAZ, O needs to be secured at 0.001% or more. The reason is that in the vicinity of the HAZ melting line, a large number of fine oxides are dispersed by appropriate addition of Ca and Mg in order to increase the HAZ toughness, and the pinning effect is strengthened to achieve γ grain refinement. If O is less than 0.001%, the number of oxides may be insufficient and a sufficient pinning effect may not be obtained.

「Ｃａ：カルシウム」０．０００３〜０．００４％、及び「Ｍｇ：マグネシウム」０．０００３〜０．００４％
Ｃａ及びＭｇは、溶鋼への添加順序を考慮しつつ、一方あるいは両方を０．０００３％以上添加することで、ＣａやＭｇを含有する１０〜５００ｎｍの酸化物や硫化物を１０００個／ｍｍ２以上確保することができる。ＣａやＭｇが０．０００３％未満の場合、大入熱溶接ＨＡＺのピン止め粒子である酸化物や硫化物の個数が不足する虞がある。また、それぞれ０．００４％超添加すると、酸化物や硫化物が粗大化してピン止め粒子の個数が不足すると同時に、破壊起点としての有害性も顕著となり、良好なＨＡＺ靭性が得られなくなる虞がある。“Ca: calcium” 0.0003 to 0.004% and “Mg: magnesium” 0.0003 to 0.004%
Ca and Mg are added in an amount of 0.0003% or more while considering the order of addition to the molten steel, so that 1000 pieces / mm 2 of oxides and sulfides of 10 to 500 nm containing Ca and Mg are added. This can be ensured. When Ca and Mg are less than 0.0003%, the number of oxides and sulfides that are pinning particles of the high heat input welding HAZ may be insufficient. Further, if each of them exceeds 0.004%, the oxides and sulfides become coarse and the number of pinning particles becomes insufficient. At the same time, the harmfulness as a starting point of fracture becomes remarkable, and there is a possibility that good HAZ toughness cannot be obtained. is there.

「Ｎｉ：ニッケル」０．０１〜１％
Ｎｉは、靭性の劣化を抑えて強度を確保するために有効である。そのためには０．０１％以上のＮｉ添加が必要である。しかしながら、Ｎｉは合金コストが非常に高いうえに、表面疵の手入工程が発生するという問題がある。従って、Ｎｉは１％以下に抑える必要がある。また、表面疵を回避するために、Ｎｉは極力低くすることが好ましいため、０．７％以下または０．５％以下に制限することがよい。"Ni: Nickel" 0.01-1%
Ni is effective for suppressing strength deterioration and ensuring strength. For that purpose, addition of 0.01% or more of Ni is necessary. However, Ni has a problem that the alloy cost is very high and a surface flawing process occurs. Therefore, Ni needs to be suppressed to 1% or less. In order to avoid surface flaws, it is preferable to reduce Ni as much as possible, so it is preferable to limit it to 0.7% or less or 0.5% or less.

「Ｃｕ：銅」０．０１〜１％、「Ｃｒ：クロム」０．０１〜１％、及び「Ｍｏ：モリブデン」０．０１〜０．５％
Ｃｕ、Ｃｒ、及びＭｏは、強度を確保するために有効であり、ともに０．０１％以上の添加量で効果を発揮する。一方、大入熱溶接ＨＡＺ靭性を劣化させる観点から、それぞれ１％、１％、及び０．５％が上限であり、それぞれ０．４％、０．３％、０．１％以下とすることが望ましい。特に、ＣｒおよびＭｏはＮｉ同様に高価な元素であり、さらにＨＡＺのＭＡ生成を助長する危険性も高いので、添加しないことが好ましい。"Cu: Copper" 0.01-1%, "Cr: Chromium" 0.01-1%, and "Mo: Molybdenum" 0.01-0.5%
Cu, Cr, and Mo are effective for securing the strength, and both exhibit an effect when added in an amount of 0.01% or more. On the other hand, from the viewpoint of degrading the high heat input welding HAZ toughness, the upper limits are 1%, 1%, and 0.5%, respectively, and should be 0.4%, 0.3%, and 0.1% or less, respectively. Is desirable. In particular, Cr and Mo are expensive elements like Ni, and further, there is a high risk of promoting MA formation of HAZ, so it is preferable not to add them.

「Ｎｂ：ニオブ」０．００３〜０．０３％
Ｎｂは、仕上圧延における未再結晶域圧延を促すために有効である。そのためには０．００３％以上のＮｂ添加が好ましい。また、Ｎｂは、大入熱溶接ＨＡＺ靭性に対しては有害である。従って、本発明では未再結晶域圧延を促すために０．０３％以下の微量のＮｂを添加してもよい。ＨＡＺ靭性の観点からは、０．０２％以下または０．０１％以下に抑えることがよい。また、仕上圧延での累積圧下量を大きく確保できる場合には、Ｎｂ無添加でも充分に母材組織が微細化して良好な脆性破壊伝播停止特性が得られるため、Ｎｂを添加しないことがＨＡＺ靭性の観点からさらに好ましい。"Nb: Niobium" 0.003-0.03%
Nb is effective for promoting non-recrystallization zone rolling in finish rolling. For that purpose, Nb addition of 0.003% or more is preferable. Nb is also harmful to high heat input welding HAZ toughness. Accordingly, in the present invention, a small amount of Nb of 0.03% or less may be added to promote non-recrystallization zone rolling. From the viewpoint of HAZ toughness, it is preferable to keep it to 0.02% or less or 0.01% or less. In addition, when the cumulative reduction amount in finish rolling can be ensured, the base metal structure can be sufficiently refined even when Nb is not added, and good brittle fracture propagation stop characteristics can be obtained. From the viewpoint of

「ＲＥＭ：希土類元素（ランタノイド系元素）」０．０００３〜０．０２％、及び「Ｚｒ：ジルコニウム」０．０００３〜０．０２％
ＲＥＭ（希土類元素）及びＺｒは、脱酸と脱硫に関与して、中心偏析部の粗大な延伸ＭｎＳの生成を抑えて硫化物を球状無害化し、母材と大入熱溶接ＨＡＺの靭性を改善する。これらの効果を発揮するためには、ＲＥＭとＺｒの下限はともに０．０００３％である。但し、これらの添加量を増やしても効果は飽和するため、経済性の観点からＲＥＭとＺｒの上限はともに０．０２％である。なお、本発明で添加するＲＥＭとは、ＬａやＣｅなどのランタノイド系元素である。"REM: rare earth element (lanthanoid element)" 0.0003 to 0.02%, and "Zr: zirconium" 0.0003 to 0.02%
REM (rare earth element) and Zr are involved in deoxidation and desulfurization, suppress the formation of coarse stretched MnS in the center segregation part, detoxify sulfides, and improve the toughness of the base metal and high heat input welding HAZ. To do. In order to exert these effects, the lower limits of REM and Zr are both 0.0003%. However, even if these addition amounts are increased, the effect is saturated, so the upper limits of REM and Zr are both 0.02% from the viewpoint of economy. The REM added in the present invention is a lanthanoid element such as La or Ce.

以上説明したように、本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法、及び脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板によれば、鋼成分を、各元素を上記範囲で含有するとともに上記各関係式を満たす成分組成とし、また、上記各製造条件とすることにより、（１）板厚５０〜８０ｍｍ、降伏強度３９０〜４６０ＭＰａ級（すなわち引張強度５１０〜５７０ＭＰａ級）の厚手高強度で、（２）アレスト性指標Ｔｋｃａ＝６０００≦−１０℃の良好な脆性破壊伝播停止特性を有し、（３）溶接入熱量≧２０ｋＪ／ｍｍでもｖＥ（−２０℃）≧４７Ｊとなる良好な大入熱溶接ＨＡＺ靭性を有し、（４）高価合金元素の低減（Ｎｉ≦０．５％等）等による低い製造コストを実現した厚手高強度鋼板が得られる。
このような、本発明による厚手高強度鋼板が大型船舶をはじめとする各種の溶接構造物に使用されることで、溶接構造物の大型化、破壊に対する高い安全性、建造における溶接の高能率化、素材である鋼材の経済性等々が同時に満たされことから、その産業上の効果は計り知れない。As described above, the method for producing a thick high-strength steel sheet having excellent brittle fracture propagation stop characteristics and high heat input heat affected zone toughness according to the present invention, and brittle fracture propagation stop characteristics and large heat input weld heat affected zone toughness. According to the thick high-strength steel sheet excellent in the above, (1) the thickness of the steel component by setting the steel component to a component composition that contains each element in the above range and satisfies the above relational expressions, and the above production conditions. Thick high strength of 50 to 80 mm, yield strength of 390 to 460 MPa class (that is, tensile strength of 510 to 570 MPa class), and (2) arrestability index T kca = 6000 ≦ −10 ° C. (3) It has good large heat input welding HAZ toughness such that vE (−20 ° C.) ≧ 47 J even when the welding heat input ≧ 20 kJ / mm, and (4) Reduction of expensive alloy elements (Ni ≦ 0.5%, etc.) ) Low manufacturing by etc. Thick high-strength steel sheet that achieves a strike can be obtained.
Such thick, high-strength steel sheets according to the present invention are used in various types of welded structures including large ships, so that the welded structures can be made larger, safe against breakage, and more efficient in welding. Because the economics of steel, the material, etc. are satisfied at the same time, the industrial effects are immeasurable.

以下、本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法、及び脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の実施例を挙げ、本発明をより具体的に説明するが、本発明は、もとより下記実施例に限定されるものではなく、前記および後記の趣旨に適合し得る範囲で適当に変更を加えて実施することも可能であり、それらはいずれも本発明の技術的範囲に含まれるものである。   Hereinafter, a method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and large heat input welding heat affected zone toughness according to the present invention, and a thick material excellent in brittle fracture propagation stopping characteristics and large heat input welding heat affected zone toughness Examples of the high-strength steel sheet will be given and the present invention will be described more specifically, but the present invention is not originally limited to the following examples, and is appropriately modified within a range that can be adapted to the purpose described above and below. It is also possible to carry out by adding these, and they are all included in the technical scope of the present invention.

［サンプル作製］
製鋼工程において溶鋼の脱酸・脱硫と化学成分を制御し、連続鋳造によって下記表１〜４に示す化学成分のスラブ（連続鋳造スラブ）を作製した。そして、下記表５〜１０に示す製造条件で、前記スラブを再加熱して厚板圧延することで板厚５０〜８０ｍｍに仕上げ、加速冷却を行い、さらに、必要に応じてオフラインでの焼戻し処理を行い、厚手鋼板のサンプルを作製した。[Sample preparation]
In the steel making process, deoxidation / desulfurization of the molten steel and chemical components were controlled, and slabs (chemical casting slabs) having chemical components shown in Tables 1 to 4 below were produced by continuous casting. Then, under the production conditions shown in Tables 5 to 10 below, the slab is reheated and rolled into a thick plate to finish the plate to a thickness of 50 to 80 mm, accelerated cooling is performed, and offline tempering treatment is performed as necessary. A thick steel plate sample was prepared.

本実施例における、本発明鋼の厚手鋼板の化学成分組成の一覧を表１，２に示すとともに、比較鋼の化学成分組成の一覧を表３，４に示す。また、本発明鋼の鋼板の製造条件の一覧を表５，６に示すとともに、比較鋼の鋼板の製造条件の一覧を表７，８に示す。また、表１，２に示す本発明鋼の「鋼No.１」の化学成分組成で、各製造条件を変化させて鋼板を製造した比較鋼の条件一覧を表９，１０に示す。   Tables 1 and 2 list the chemical composition of the thick steel plate of the present invention steel in this example, and Tables 3 and 4 list the chemical composition of the comparative steel. In addition, Tables 5 and 6 show a list of manufacturing conditions for steel plates of the present invention steel, and Tables 7 and 8 show a list of manufacturing conditions for steel plates for comparative steel. In addition, Tables 9 and 10 show a list of conditions for comparative steels manufactured by changing the respective production conditions with the chemical composition of “steel No. 1” of the steels of the present invention shown in Tables 1 and 2.

なお、表２，４において、Ｃｅｑ、Ａ式、Ｂ式、Ｃ式、Ｄ式、及びＡｒ３は、以下のように規定される。
Ｃｅｑ＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５
Ａ式＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ
Ｂ式＝Ｎ−０．２９Ｔｉ
Ｃ式＝Ｔｉ−２（Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ）
Ｄ式＝Ｎ−０．２９〔Ｔｉ−２（Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ）〕
Ａｒ３＝９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ （スラブ）
また、有効Ｂ量は、以下のように規定される。
（ｉ）Ａ式の値＜０の場合
（ａ）Ｂ式の値＞０の場合、有効Ｂ量＝Ｂ−０．７７（Ｎ−０．２９Ｔｉ)
（ｂ）Ｂ式の値≦０の場合、有効Ｂ量＝Ｂ
（ｉｉ）Ａ式の値≧０の場合
Ｃ式の値≧０．００５
（ａ）Ｄ式の値＞０の場合、有効Ｂ量＝Ｂ−０．７７｛Ｎ−０．２９〔Ｔｉ−２（Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ）〕｝
（ｂ）Ｄ式の値≦０の場合、有効Ｂ量＝ＢIn Tables 2 and 4, Ceq, A formula, B formula, C formula, D formula, and Ar 3 are defined as follows.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15
Formula A = O-0.4Ca-0.66Mg-0.17REM-0.35Zr-0.89Al
Formula B = N-0.29Ti
Formula C = Ti-2 (O-0.4Ca-0.66Mg-0.17REM-0.35Zr-0.89Al)
Formula D = N-0.29 [Ti-2 (O-0.4Ca-0.66Mg-0.17REM-0.35Zr-0.89Al)]
Ar 3 = 910-310C-80Mn-20Cu-55Ni-80Mo (slab)
Further, the effective B amount is defined as follows.
(I) When the value of the formula A <0 (a) When the value of the formula B> 0, the effective B amount = B−0.77 (N−0.29Ti)
(B) When the value of the formula B ≦ 0, the effective B amount = B
(Ii) When the value of the formula A ≧ 0 The value of the formula C ≧ 0.005
(A) When the value of the formula D> 0, the effective B amount = B−0.77 {N−0.29 [Ti-2 (O−0.4Ca−0.66Mg−0.17REM−0.35Zr− 0.89Al)]}
(B) When the value of the formula D ≦ 0, the effective B amount = B

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

Figure 2009072559
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Figure 2009072559
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Figure 2009072559
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Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

［評価試験］
上記方法によって作製した厚手鋼板のサンプルについて、以下のような評価試験を行った。
母材の引張特性及びシャルピー衝撃特性については、厚手鋼板サンプルの板厚１／２部−圧延長手（Ｌ）方向から試験片を採取して測定して評価した。
母材の脆性破壊伝播停止特性については、全厚試験体を温度勾配型ＥＳＳＯ試験（ＷＥＳ ３００３準拠）によって破壊し、アレスト性指標Ｔｋｃａ＝６０００を求めて評価した。
継手のＨＡＺ靭性については、突合せ開先をエレクトロガス溶接（ＥＧＷ）によって１パス溶接し、板厚１／２部の溶融線から１ｍｍ離れたＨＡＺにノッチを入れて調べた。この際、−２０℃で３本のシャルピー衝撃試験を行い、平均の吸収エネルギー値を評価した。また、参考として、−４０℃における特性も調べた。[Evaluation test]
The following evaluation tests were performed on samples of thick steel plates produced by the above method.
The tensile properties and Charpy impact properties of the base material were evaluated by collecting and measuring test pieces from the plate thickness 1/2 part of the thick steel plate sample-rolling longitudinal (L) direction.
The brittle fracture propagation stop characteristics of the base material were evaluated by destroying the full thickness specimen by a temperature gradient type ESSO test (based on WES 3003) and obtaining an arrestability index T kca = 6000 .
Regarding the HAZ toughness of the joint, the butt groove was welded by one pass by electrogas welding (EGW), and a notch was made in the HAZ 1 mm away from the melt line having a thickness of 1/2 part. At this time, three Charpy impact tests were performed at −20 ° C., and the average absorbed energy value was evaluated. For reference, the characteristics at −40 ° C. were also examined.

厚手鋼板と溶接継手の機械的性質について、表５，６に示す製造条件で製造した本発明鋼の機械的性質一覧を表１１に示し、また、表７，８に示す製造条件で製造した比較鋼の機械的性質一覧を表１２に示す。
本発明鋼の「鋼No.１」の化学成分組成で、製造条件を表９，１０に示す条件で変化させて製造した比較鋼の、厚手鋼板と溶接継手の機械的性質一覧を表１３に示す。About the mechanical properties of the thick steel plate and the welded joint, a list of mechanical properties of the steels of the present invention manufactured under the manufacturing conditions shown in Tables 5 and 6 is shown in Table 11, and a comparison manufactured under the manufacturing conditions shown in Tables 7 and 8 Table 12 shows a list of mechanical properties of the steel.
Table 13 shows a list of mechanical properties of thick steel plates and welded joints of comparative steels manufactured by changing the manufacturing conditions under the conditions shown in Tables 9 and 10 with the chemical composition of steel No. 1 of the present invention steel. Show.

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

Figure 2009072559
Figure 2009072559

［評価結果］
表１，２に示す鋼No.１〜１６は本発明鋼であり、鋼の化学成分を適正化し、ＴＭＣＰにおける低温加熱と低温圧延を徹底することにより、厚手であるのにも関わらず、表１１に示すように、３９０〜４６０ＭＰａ級の降伏強度と５１０〜５６０ＭＰａ級の引張強度、及び、−１０℃未満の良好な脆性破壊伝播停止特性Ｔｋｃａ＝６０００を満足し、さらに、大入熱溶接であるのにも関わらず、−２０℃において良好なＨＡＺ靭性が、Ｎｉ添加量を１％以下に抑えながら、同時に満足できていることがわかる。[Evaluation results]
Steel Nos. 1 to 16 shown in Tables 1 and 2 are steels of the present invention. By optimizing the chemical composition of the steel and thoroughly implementing low-temperature heating and low-temperature rolling in TMCP, 11 satisfies 390 to 460 MPa class yield strength, 510 to 560 MPa class tensile strength, and favorable brittle fracture propagation stop property T kca = 6000 below -10 ° C. Nevertheless, it can be seen that good HAZ toughness at −20 ° C. is satisfied at the same time while suppressing the Ni addition amount to 1% or less.

一方、表３，４に示す比較鋼No.１７〜３６は、鋼の化学成分が適正でなく、また、表９，１０に示す比較鋼１Ａ〜１Ｉは鋼板製造条件が適正でないため、表１２，１３に示すように、降伏強度、引張強度、Ｔｋｃａ＝６０００、及び大入熱溶接ＨＡＺ靭性のいずれかが劣り、本発明の厚手高強度鋼板のように、これら複数の要求特性を同時に満足することができないことがわかる。On the other hand, Comparative Steel Nos. 17 to 36 shown in Tables 3 and 4 are not appropriate in chemical composition of steel, and Comparative Steels 1A to 1I shown in Tables 9 and 10 are not appropriate in steel plate manufacturing conditions. 13, any one of yield strength, tensile strength, T kca = 6000 , and high heat input weld HAZ toughness is inferior, and simultaneously satisfies these multiple required properties as in the thick high strength steel sheet of the present invention. You can't do it.

鋼No.１７はＣとＣｅｑが低いため、また、鋼No.２０はＭｎが低いため、焼入性が不足して降伏強度や引張強度が劣っている。
鋼No.１８はＣが高いため、鋼No.１９はＳｉが高いため、鋼No.２１はＭｎが高いため、また、鋼No.２２はＢが低いため、それぞれ大入熱溶接ＨＡＺの靭性が劣っている。
鋼No.２３はＶが低いため、板厚が同じでＣｅｑが低い鋼No.１より強度が低く、また鋼No.１よりＣｅｑが高いにもかかわらず、鋼No.１が満たす４６０ＭＰａ級降伏強度と５７０ＭＰａ級引張強度を満たすことができない。さらに、大入熱溶接ＨＡＺの靭性が劣っている。
鋼No.２４はＶが高いため、板厚とＣｅｑが同じ鋼No.１１よりも強度が大幅に高いが、大入熱溶接ＨＡＺの靭性が劣っている。
鋼No.２５、２６、２７、３０、３１、３４、及び３５はＣｅｑと板厚が同じであり、また表７，８のＴＭＣＰ条件も同一であるが、有効Ｂ量が８〜１０ｐｐｍ存在するため、降伏強度は４４０〜６００ＭＰａであり、引張強度は５５０〜７００ＭＰａであり、強度の変動が大きい。さらに、大入熱溶接ＨＡＺの靭性が劣っている。Steel No. 17 has low C and Ceq, and Steel No. 20 has low Mn. Therefore, hardenability is insufficient and yield strength and tensile strength are inferior.
Steel No. 18 is high in C, Steel No. 19 is high in Si, Steel No. 21 is high in Mn, and Steel No. 22 is low in B, so the toughness of high heat input welding HAZ respectively. Is inferior.
Steel No. 23 has a low V, so the strength is lower than Steel No. 1 with the same thickness and low Ceq, and the 460 MPa class yield that Steel No. 1 satisfies despite the higher Ceq than Steel No. 1. The strength and the 570 MPa class tensile strength cannot be satisfied. Furthermore, the toughness of the high heat input weld HAZ is inferior.
Since steel No. 24 has a high V, its strength is significantly higher than steel No. 11 having the same thickness and Ceq, but the toughness of high heat input welding HAZ is inferior.
Steel Nos. 25, 26, 27, 30, 31, 34, and 35 have the same plate thickness as Ceq, and the TMCP conditions in Tables 7 and 8 are the same, but the effective B amount is 8 to 10 ppm. Therefore, the yield strength is 440 to 600 MPa, the tensile strength is 550 to 700 MPa, and the strength varies greatly. Furthermore, the toughness of the high heat input weld HAZ is inferior.

鋼No.２８はＰが高いため、また鋼No.２９はＳが高いために、それぞれ母材靭性と大入熱ＨＡＺの靭性が劣っている。
鋼No.３１はＡｌが低いためにＯが高くなり、鋼No.３２はＡｌが高いためにアルミナクラスターが生成し、ともに粗大な有害酸化物が増えて母材と大入熱ＨＡＺの靭性が劣っている。
鋼No.３３はＴｉが低いため、また鋼No.３５はＮが低いために、それぞれＴｉＮの生成が不十分で母材とＨＡＺの結晶粒が十分に微細化されず、母材靭性、アレスト性、及び大入熱ＨＡＺ靭性が劣っている。
鋼No.３４はＴｉが高いため、また鋼No.３６はＮが高いため、それぞれＴｉＣ脆化や固溶Ｂ脆化によって母材靭性と大入熱ＨＡＺ靭性が劣っている。Since steel No. 28 has high P and steel No. 29 has high S, the toughness of the base metal and the high heat input HAZ are inferior, respectively.
Steel No. 31 has high Al due to low Al, and Steel No. 32 has high Al, so alumina clusters are formed, both of which increase the amount of coarse toxic oxides and increase the toughness of the base metal and high heat input HAZ. Inferior.
Since steel No. 33 has low Ti and steel No. 35 has low N, the formation of TiN is insufficient, and the base metal and HAZ crystal grains are not sufficiently refined. And high heat input HAZ toughness are inferior.
Steel No. 34 is high in Ti and Steel No. 36 is high in N. Therefore, the base metal toughness and the high heat input HAZ toughness are inferior due to TiC embrittlement and solute B embrittlement, respectively.

鋼No.１Ａは、スラブ再加熱の開始温度が高いため、また鋼No.１Ｂは加熱温度が高いために、それぞれ加熱時のγ粒が粗大化して脆性破壊伝播停止特性Ｔｋｃａ＝６０００が劣っている。
鋼No.１Ｃは、加熱温度が低すぎるためにＶ炭窒化物の溶体化が不十分となり、析出強化を担うＶ炭化物が不足して母材強度が低下しているため、鋼No.１に比べて降伏強度と引張強度がともに２０ＭＰａ低くなり、０．０２％のＶを添加した強度のメリットを享受できていない。さらに、粗圧延の終了温度が低すぎるために再結晶粒が十分に整細粒化されず、Ｔｋｃａ＝６０００が劣っている。
鋼No.１Ｄは、粗圧延の終了温度が低すぎるため、また、鋼No.１Ｅは粗圧延の累積圧下量が少ないために再結晶粒が十分に整細粒化されず、それぞれＴｋｃａ＝６０００が劣っている。
鋼No.１Ｆと鋼No.１Ｇは、仕上圧延の開始温度と終了温度が高すぎて上記式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝を満足しないため、母材の結晶粒径の微細化が不十分であり、Ｔｋｃａ＝６０００が劣っている。
鋼No.１Ｈは、仕上圧延の累積圧下量が少ないため、母材の結晶粒径の微細化が不十分であり、Ｔｋｃａ＝６０００が劣っている。
鋼No.１Ｉは、加速冷却の停止温度が高いため、板厚内部の変態強化と結晶粒径微細化が不十分となり、引張強度とＴｋｃａ＝６０００が劣っている。Steel No. 1A has a high starting temperature for slab reheating, and Steel No. 1B has a high heating temperature, so that the γ grains during heating are coarsened and the brittle fracture propagation stop property T kca = 6000 is inferior. ing.
In Steel No. 1C, since the heating temperature is too low, solution of V carbonitride is insufficient, and V carbide responsible for precipitation strengthening is insufficient and the base metal strength is reduced. In comparison, the yield strength and the tensile strength are both reduced by 20 MPa, and the merit of strength obtained by adding 0.02% V cannot be enjoyed. Furthermore, since the end temperature of the rough rolling is too low, the recrystallized grains are not sufficiently refined, and T kca = 6000 is inferior.
In Steel No. 1D, the end temperature of rough rolling is too low, and in Steel No. 1E, the cumulative reduction amount of rough rolling is small, so the recrystallized grains are not sufficiently refined, and T kca = 6000 is inferior.
Steel No. 1F and Steel No. 1G are not suitable for the above formula {−0.5 × (slab heating temperature (° C.)) + 1325} because the finish rolling start temperature and finish temperature are too high. The particle size is not sufficiently refined, and T kca = 6000 is inferior.
Steel No. 1H has a small amount of cumulative reduction in finish rolling, so that the crystal grain size of the base material is not sufficiently refined, and T kca = 6000 is inferior.
Steel No. 1I has a high accelerated cooling stop temperature, so that the transformation strengthening inside the plate thickness and the refinement of crystal grain size are insufficient, and the tensile strength and T kca = 6000 are inferior.

以上説明した実施例の結果より、本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板が、（１）板厚５０〜８０ｍｍ、降伏強度３９０〜４６０ＭＰａ級（すなわち引張強度５１０〜５７０ＭＰａ級）の厚手高強度で、（２）アレスト性指標Ｔｋｃａ＝６０００≦−１０℃の良好な脆性破壊伝播停止特性を有し、（３）溶接入熱量≧２０ｋＪ／ｍｍでもｖＥ（−２０℃）≧４７Ｊとなる良好な大入熱溶接ＨＡＺ靭性を有し、（４）高価合金元素の低減（Ｎｉ≦１％等）等による低い製造コストを実現できることが明らかである。From the results of the examples described above, a thick high-strength steel plate excellent in brittle fracture propagation stop characteristics and high heat input welding heat-affected zone toughness of the present invention is (1) plate thickness 50-80 mm, yield strength 390-460 MPa class. (2) Arrestability index T kca = 6000 ≦ −10 ° C. Good brittle fracture propagation stop property (3) Weld heat input ≧ 20 kJ / It is clear that even with mm, it has good high heat input HAZ toughness satisfying vE (−20 ° C.) ≧ 47 J, and (4) low manufacturing cost can be realized by reducing expensive alloy elements (Ni ≦ 1%, etc.). is there.

本発明の厚手高強度鋼板は、大型船舶をはじめとする各種の溶接構造物に使用されることで、溶接構造物の大型化、破壊に対する高い安全性、建造における溶接の高能率化、及び素材である鋼材の経済性等が同時に満たされる。このため、本発明に係る厚手高強度鋼板は、大型コンテナ船等の船舶用途や、建築、橋梁、タンク及び海洋構造物等のその他の溶接構造物に適用できる。   The thick high-strength steel sheet of the present invention is used for various welded structures including large ships, so that the welded structures are enlarged, the safety against destruction, the efficiency of welding in construction, and the material The economics etc. of steel materials which are are satisfy | filled simultaneously. For this reason, the thick high-strength steel sheet according to the present invention can be applied to ship applications such as large container ships and other welded structures such as buildings, bridges, tanks and offshore structures.

上記問題を解決するための本発明の要旨は以下のとおりである。
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部が鉄および不可避的不純物からなる連続鋳造スラブをＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱し、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を前記連続鋳造スラブに行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とし、次いで、加速冷却を適用して前記圧延原板を５００℃以下まで冷却して鋼板とする。前記連続鋳造スラブは、有効Ｂ量：Ｂｅｆ（％）の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たす。
ここで、残存酸素量ＯＴｉ（％）を、下記式（１）で表される量としたとき、有効Ｂ量：Ｂｅｆ（％）は、下記式（２）で表される。また炭素当量Ｃｅｑ（％）は、下記式（３）で表され、Ａｒ３は、下記式（４）で表される。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ・・・（１）｛但し、式（１）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）｛但し、式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（３）
Ａｒ３（℃）＝（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ） ・・・（４）
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法では、前記加速冷却の後、さらに、３５０〜７００℃で５〜６０分の焼戻し熱処理を施してもよい。
前記連続鋳造スラブの前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、前記連続鋳造スラブは、さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有してもよい。
前記連続鋳造スラブは、さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有してもよい。
前記連続鋳造スラブは、さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有してもよい。 The gist of the present invention for solving the above problems is as follows.
The manufacturing method of the thick high-strength steel sheet excellent in brittle fracture propagation stop characteristics and high heat input welding heat-affected zone toughness according to the present invention is mass%, C: 0.05 to 0.12%, Si: 0.3%. Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al: 0 .001~0.1%, Ti: 0.005~0.02%, N: 0.002~0.01%, and O: to 0.004% or less, the balance being iron and unavoidable impurities The continuous cast slab is cooled to Ar 3 −200 ° C. or lower and then reheated to 950 to 1100 ° C., and then subjected to rough rolling at 900 ° C. or higher and a cumulative reduction amount of 30% or higher on the continuous cast slab, Next, finish rolling with a cumulative reduction amount of 50% or more at 700 ° C. or higher is applied to the finish rolling start temperature. And the finish rolling finish temperature are both set to the temperature represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.) or less to obtain a rolled original plate, and then accelerated. The rolling original sheet is cooled to 500 ° C. or less by applying cooling to obtain a steel sheet. The continuously cast slab, effective B amount: calculated value Bef (%) is not more than 0%, the carbon equivalent Ceq satisfy the range of from 0.32 to 0.42%.
Here, no residual oxygen quantity O Ti (%), when the amount represented by the following formula (1), effective B amount: Bef (%) is represented by the following formula (2). The carbon equivalent Ceq (%) is represented by the following formula (3), and Ar 3 is represented by the following formula (4).
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1) {However, in the formula (1), the component elements treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (2) {However, in the formula (2), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (3)
Ar 3 (° C.) = (910-310C-80Mn-20Cu-55Ni-80Mo) (4)
In the method for producing a thick high-strength steel sheet having excellent brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention, after the accelerated cooling, tempering heat treatment is further performed at 350 to 700 ° C. for 5 to 60 minutes. You may give it.
The content of S in the continuous cast slab is 0.0005 to 0.005%, and the content of O is 0.001 to 0.004%. Thus, one or two of Ca: 0.0003 to 0.004% and Mg: 0.0003 to 0.004% may be contained.
The continuous cast slab further comprises, in mass%, Ni: 0.01 to 1%, Cu: 0.01 to 1%, Cr: 0.01 to 1%, Mo: 0.01 to 0.5%, And Nb: You may contain 1 type (s) or 2 or more types in 0.003-0.03%.
The continuous cast slab may further contain one or two of REM: 0.0003 to 0.02% and Zr: 0.0003 to 0.02% by mass%.

本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部が鉄および不可避的不純物からなり、残存酸素量Ｏ Ｔｉ （％）を、下記式（５）で表される量としたとき、下記式（６）で表される有効Ｂ量：Ｂｅｆ（％）の算出値が０％以下であり、さらに、下記式（７）で表される炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たし、板厚が５０〜８０ｍｍであり、降伏強度が３９０〜４６０ＭＰａ級であり、脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００が−１０℃以下であり、溶接入熱量が２０ｋＪ／ｍｍ以上の大入熱溶接部のＨＡＺ靭性の指標であるシャルピー衝撃吸収エネルギーｖＥ（−２０℃）が４７Ｊ以上である。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ・・・（５）｛但し、式（５）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（６）｛但し、式（６）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（７）
本発明の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板では、前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有してもよい。
さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有してもよい。
さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有してもよい。 The thick high-strength steel sheet having excellent brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention is mass%, C: 0.05 to 0.12%, Si: 0.3% or less, Mn : 1-2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003-0.003%, V: 0.01-0.15%, Al: 0.001- 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, and O: 0.004% or less, with the balance being iron and unavoidable impurities , presence of oxygen quantity O Ti a (%), when the amount represented by the following formula (5), effective B amount represented Ru by the following formula (6): calculated value of the Bef (%) is not more than 0% Furthermore, the carbon equivalent Ceq represented by the following formula (7) satisfies the range of 0.32 to 0.42%, the plate thickness is 50 to 80 mm, and the yield strength is 3 A 0~460MPa grade, brittle fracture propagation stop characteristics Kca is at a temperature T kca = 6000 to be 6000 N / mm 1.5 is -10 ° C. or less, the welding heat input of high heat input welds than 20 kJ / mm Charpy impact absorption energy vE (−20 ° C.), which is an index of HAZ toughness, is 47 J or more.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (5) {However, in the formula (5), component elements treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (6) {However, in the formula (6), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (7)
In the thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness of the present invention, the content of S is 0.0005 to 0.005%, and the content of O is 0.001 to 0.004%, and further, by mass%, containing one or two of Ca: 0.0003 to 0.004% and Mg: 0.0003 to 0.004% Also good.
Furthermore, by mass%, Ni: 0.01-1%, Cu: 0.01-1%, Cr: 0.01-1%, Mo: 0.01-0.5%, and Nb: 0.003 You may contain 1 type, or 2 or more types in -0.03%.
Furthermore, you may contain 1 type or 2 types in REM: 0.0003-0.02% and Zr: 0.0003-0.02% by the mass%.

＜鋼板製造条件（製造方法）＞
船舶等の溶接構造物に使用される鋼板においては、（１）大きな板厚での高い強度、（２）良好な脆性破壊伝播停止特性、（３）良好な大入熱溶接ＨＡＺ靭性、及び（４）低い製造コスト等のニーズが高まっている。
このようなニーズに対し、本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法は、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ ：０．００４％以下を含有し、残部が鉄および不可避的不純物からなる連続鋳造スラブを、連続鋳造後にＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱する工程と、次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を連続鋳造スラブに行い、次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とする工程と、次いで、加速冷却を適用して圧延原板を５００℃以下まで冷却して鋼板とする工程とを有する。前記連続鋳造スラブは、有効Ｂ量：Ｂｅｆ（％）の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たす。
ここで、残存酸素量ＯＴｉ（％）を、下記式（１）で表される量としたとき、有効Ｂ量：Ｂｅｆ（％）は、下記式（２）で表される。また炭素当量Ｃｅｑ（％）は、下記式（３）で表され、Ａｒ３は、下記式（４）で表される。
また、スラブ加熱温度とは、連続鋳造スラブを再加熱する際の温度（再加熱温度）である。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（１）｛但し、式（１）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）｛但し、式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（３）
Ａｒ３（℃）＝（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ） ・・・（４）
なお、本明細書において、式中の元素記号は、連続鋳造スラブ又は厚手高強度鋼板中のその元素の含有量（質量％）を示す。
また、本発明において連続鋳造スラブの製造方法は特に限定されない。例えば、高炉、転炉や電炉等による溶製に引き続き、各種の２次精練で目的の成分含有量になるように成分調整を行い、次いで通常の連続鋳造によって製造される。 <Steel plate manufacturing conditions (manufacturing method)>
In steel plates used for welded structures such as ships, (1) high strength at large plate thickness, (2) good brittle fracture propagation stop properties, (3) good high heat input weld HAZ toughness, and ( 4) Needs such as low manufacturing costs are increasing.
For such needs, the method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness according to the present invention is in mass%, and C: 0.05 to 0.12. %, Si: 0.3% or less, Mn: 1-2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003-0.003%, V: 0.01- 0.15%, Al: 0.001-0.1%, Ti: 0.005-0.02%, N: 0.002-0.01%, and O: 0.004% or less, after the remainder of the continuous cast slab comprising iron and unavoidable impurities, and cooled to Ar 3 -200 ° C. or less after the continuous casting, the steps of reheating to 950 to 1,100 ° C., then the cumulative reduction ratio at 900 ° C. or higher 30 % Is performed on a continuous cast slab, and then the cumulative reduction amount is 700 ° C. or more. In the finish rolling that is 0% or more, the finish rolling start temperature and the finish rolling end temperature are both below the temperature represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325} (° C.) And a step of forming a rolled original sheet under the conditions described above, and then a step of applying accelerated cooling to cool the rolled original sheet to 500 ° C. or lower to obtain a steel sheet. The continuously cast slab, effective B amount: calculated value Bef (%) is not more than 0%, the carbon equivalent Ceq satisfy the range of from 0.32 to 0.42%.
Here, no residual oxygen quantity O Ti (%), when the amount represented by the following formula (1), effective B amount: Bef (%) is represented by the following formula (2). The carbon equivalent Ceq (%) is represented by the following formula (3), and Ar 3 is represented by the following formula (4).
The slab heating temperature is a temperature (reheating temperature) when reheating the continuously cast slab.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1) {However, in the formula (1), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (2) {However, in the formula (2), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (3)
Ar 3 (° C.) = (910-310C-80Mn-20Cu-55Ni-80Mo) (4)
In addition, in this specification, the element symbol in a type | formula shows content (mass%) of the element in a continuous casting slab or a thick high-strength steel plate.
In the present invention, the method for producing the continuous cast slab is not particularly limited. For example, subsequent to melting by a blast furnace, a converter, an electric furnace, etc., the components are adjusted so that the desired component content is obtained by various secondary scouring, and then manufactured by ordinary continuous casting.

＜化学成分組成（厚手高強度鋼板）＞
本発明に係る脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板は、上述したような、（１）大きな板厚での高い強度、（２）良好な脆性破壊伝播停止特性、（３）良好な大入熱溶接ＨＡＺ靭性、及び（４）低い製造コスト等のニーズを満足するため、質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部が鉄および不可避的不純物からなり、残存酸素量Ｏ Ｔｉ （％）を、下記式（５）で表される量としたとき、下記式（６）で表される有効Ｂ量：Ｂｅｆ（％）の算出値が０％以下であり、さらに、下記式（７）で表される炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たし、板厚が５０〜８０ｍｍであり、降伏強度が３９０〜４６０ＭＰａ級で、引張強度が５１０〜５７０ＭＰａ級であり、脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００が−１０℃以下であり、溶接入熱量が２０ｋＪ／ｍｍ以上の大入熱溶接部のＨＡＺ靭性の指標であるシャルピー衝撃吸収エネルギーｖＥ（−２０℃）が４７Ｊ以上とされている。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（５）｛但し、式（５）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（６）｛但し、式（６）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（７）
なお、上記式（５）〜（７）の各式において、式（５）は上記式（１）と共通の式であり、また、式（６）は上記式（２）、式（７）は上記式（３）とそれぞれ共通の式である。 <Chemical component composition (thick high-strength steel plate)>
The thick high-strength steel sheet excellent in brittle fracture propagation stop characteristics and high heat input welding heat-affected zone toughness according to the present invention is as described above, (1) high strength at a large thickness, and (2) good brittle fracture. In order to satisfy the needs such as propagation stop characteristics, (3) good high heat input HAZ toughness, and (4) low production cost, C: 0.05 to 0.12%, Si: 0.00. 3% or less, Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0003 to 0.003%, V: 0.01 to 0.15%, Al : 0.001 to 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01%, and O: 0.004% or less, the balance being iron and inevitable consists impurities, no residual oxygen quantity O Ti (%), when the amount represented by the following formula (5), you express the following formula (6) Yes B amount: Bef (%) and the calculated value of 0% or less, further, satisfies the range carbon equivalent Ceq of 0.32 to 0.42% of the following formula (7), the plate thickness is 50 80 mm, yield strength is 390 to 460 MPa class, tensile strength is 510 to 570 MPa class, brittle fracture propagation stop characteristic Kca is 6000 N / mm 1.5 , temperature T kca = 6000 is −10 ° C. or less The Charpy impact absorption energy vE (−20 ° C.), which is an index of the HAZ toughness of a large heat input weld having a welding heat input of 20 kJ / mm or more, is 47 J or more.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (5) {However, in the formula (5), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (6) {However, in the formula (6), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (7)
In each of the above formulas (5) to (7), formula (5) is a common formula with formula (1), and formula (6) is formula (2) and formula (7). Are common to the above equation (3).

Claims (9)

質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含む連続鋳造スラブをＡｒ３−２００℃以下まで冷却した後、９５０〜１１００℃に再加熱し、
次いで、９００℃以上で累積圧下量が３０％以上である粗圧延を前記連続鋳造スラブに行い、
次いで、７００℃以上で累積圧下量が５０％以上である仕上圧延を、仕上圧延開始温度および仕上圧延終了温度が、ともに、次式｛−０．５×（スラブ加熱温度（℃））＋１３２５｝（℃）で表される温度以下とされた条件で行って圧延原板とし、
次いで、加速冷却を適用して前記圧延原板を５００℃以下まで冷却して鋼板とし、
前記連続鋳造スラブは、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たすことを特徴とする、脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法。
ここで、強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量ＯＴｉ（％）を、下記式（１）で表される量としたとき、有効Ｂ量：Ｂｅｆ（％）は、下記式（２）で表される。また炭素当量Ｃｅｑ（％）は、下記式（３）で表され、Ａｒ３は、下記式（４）で表される。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（１）｛但し、式（１）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（２）｛但し、式（２）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（３）
Ａｒ３（℃）＝（９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−５５Ｎｉ−８０Ｍｏ） ・・・（４）In terms of mass%, C: 0.05 to 0.12%, Si: 0.3% or less, Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.0. 0003-0.003%, V: 0.01-0.15%, Al: 0.001-0.1%, Ti: 0.005-0.02%, N: 0.002-0.01% And, after cooling a continuous cast slab containing O: 0.004% or less and containing iron and inevitable impurities as the balance to Ar 3 −200 ° C. or less, reheating to 950 to 1100 ° C.,
Next, the continuous cast slab is subjected to rough rolling with a cumulative reduction of 30% or more at 900 ° C. or higher,
Next, finish rolling in which the cumulative reduction amount is 50% or more at 700 ° C. or higher has a finish rolling start temperature and finish rolling end temperature both represented by the following formula {−0.5 × (slab heating temperature (° C.)) + 1325}. (1) It is performed under the condition of the temperature represented by (° C.) to make a rolled original sheet,
Next, accelerated cooling is applied to cool the rolled original sheet to 500 ° C. or less to form a steel sheet,
In the continuous cast slab, the calculated amount of B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation is 0% or less, and the carbon equivalent Ceq is 0.32 to 0.42%. A method for producing a thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input welding heat-affected zone toughness, characterized by satisfying the range.
Here, when the residual oxygen amount O Ti (%) that remains after deoxidation by the strong deoxidation element and can be deoxidized by Ti, which is the weak deoxidation element, is effective when the amount is represented by the following formula (1) B amount: Bef (%) is represented by the following formula (2). The carbon equivalent Ceq (%) is represented by the following formula (3), and Ar 3 is represented by the following formula (4).
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (1) {However, in the formula (1), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (2) {However, in the formula (2), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (3)
Ar 3 (° C.) = (910-310C-80Mn-20Cu-55Ni-80Mo) (4) 前記加速冷却の後、さらに、３５０〜７００℃で５〜６０分の焼戻し熱処理を施すことを特徴とする、請求項１に記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法。   The accelerating cooling is followed by tempering heat treatment at 350 to 700 ° C for 5 to 60 minutes, and is excellent in brittle fracture propagation stopping characteristics and high heat input welding heat affected zone toughness. A manufacturing method for thick and high strength steel sheets 前記連続鋳造スラブの前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、
前記連続鋳造スラブは、さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有することを特徴とする、請求項１又は２に記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法。The content of S in the continuous cast slab is 0.0005 to 0.005%, and the content of O is 0.001 to 0.004%,
The continuous cast slab is further characterized by containing one or two of Ca: 0.0003 to 0.004% and Mg: 0.0003 to 0.004% by mass%. The manufacturing method of the thick high-strength steel plate excellent in the brittle fracture propagation stop characteristic of Claim 1 or 2, and the high heat input welding heat affected zone toughness. 前記連続鋳造スラブは、さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有することを特徴とする、請求項１〜３のいずれかに記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法。   The continuous cast slab further comprises, in mass%, Ni: 0.01 to 1%, Cu: 0.01 to 1%, Cr: 0.01 to 1%, Mo: 0.01 to 0.5%, And Nb: 0.003 to 0.03% of one type or two or more types are included, and the brittle fracture propagation stop property and high heat input welding according to any one of claims 1 to 3 A method for producing thick, high-strength steel sheets with excellent heat-affected zone toughness. 前記連続鋳造スラブは、さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有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破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板の製造方法。   The continuous cast slab is further characterized by containing one or two of REM: 0.0003 to 0.02% and Zr: 0.0003 to 0.02% by mass%. The manufacturing method of the thick high strength steel plate excellent in the brittle fracture propagation stop characteristic in any one of Claims 1-4, and the high heat input welding heat affected zone toughness. 質量％で、Ｃ：０．０５〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１〜２％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ｂ：０．０００３〜０．００３％、Ｖ：０．０１〜０．１５％、Ａｌ：０．００１〜０．１％、Ｔｉ：０．００５〜０．０２％、Ｎ：０．００２〜０．０１％、及びＯ：０．００４％以下を含有し、残部として鉄および不可避的不純物を含み、
強脱酸元素による脱酸後に残存し弱脱酸元素であるＴｉにより脱酸され得る残存酸素量を、下記式（５）で表される量としたとき、下記式（６）で表される、変態前のオーステナイト素地に固溶するＢ量｛有効Ｂ量：Ｂｅｆ（％）｝の算出値が０％以下であり、
さらに、下記式（７）で表される炭素当量Ｃｅｑが０．３２〜０．４２％の範囲を満たし、
板厚が５０〜８０ｍｍであり、降伏強度が３９０〜４６０ＭＰａ級であり、
脆性破壊伝播停止特性Ｋｃａが６０００Ｎ／ｍｍ１．５となる温度Ｔｋｃａ＝６０００が−１０℃以下であり、
溶接入熱量が２０ｋＪ／ｍｍ以上の大入熱溶接部のＨＡＺ靭性の指標であるシャルピー衝撃吸収エネルギーｖＥ（−２０℃）が４７Ｊ以上であることを特徴とする、脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板。
ＯＴｉ（％）＝Ｏ−０．４Ｃａ−０．６６Ｍｇ−０．１７ＲＥＭ−０．３５Ｚｒ−０．８９Ａｌ ・・・（５）｛但し、式（５）において、不可避的不純物扱いの成分元素も計算に含める｝
Ｂｅｆ（％）＝Ｂ−０．７７｛Ｎ−０．２９（Ｔｉ−２ＯＴｉ）｝ ・・・（６）｛但し、式（６）において、ＯＴｉ≦０のとき、ＯＴｉ＝０とする。また、ＯＴｉ＞０のときは、Ｔｉ−２ＯＴｉ≧０．００５（％）を満たすものとする。さらに、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）≦０（但し、ＯＴｉ≦０のとき、ＯＴｉ＝０）のときは、Ｎ−０．２９（Ｔｉ−２ＯＴｉ）＝０とする。｝
Ｃｅｑ（％）＝Ｃ＋Ｍｎ／６＋（Ｃｒ＋Ｍｏ＋Ｖ）／５＋（Ｎｉ＋Ｃｕ）／１５ ・・・（７）By mass%, C: 0.05 to 0.12%, Si: 0.3% or less, Mn: 1 to 2%, P: 0.015% or less, S: 0.005% or less, B: 0.00. 0003 to 0.003%, V: 0.01 to 0.15%, Al: 0.001 to 0.1%, Ti: 0.005 to 0.02%, N: 0.002 to 0.01% , And O: 0.004% or less, with iron and inevitable impurities as the balance,
When the amount of residual oxygen that remains after deoxidation by the strong deoxidation element and can be deoxidized by Ti, which is the weak deoxidation element, is represented by the following formula (5), it is represented by the following formula (6). The calculated value of B amount {effective B amount: Bef (%)} dissolved in the austenite substrate before transformation is 0% or less,
Furthermore, the carbon equivalent Ceq represented by the following formula (7) satisfies the range of 0.32 to 0.42%,
The plate thickness is 50 to 80 mm, the yield strength is 390 to 460 MPa class,
The temperature T kca = 6000 at which the brittle fracture propagation stop characteristic Kca is 6000 N / mm 1.5 is −10 ° C. or less,
Charpy impact absorption energy vE (−20 ° C.), which is an index of HAZ toughness of high heat input welds with a welding heat input of 20 kJ / mm or more, is 47 J or more, and a brittle fracture propagation stop characteristic and large input Thick high-strength steel sheet with excellent heat-welding heat affected zone toughness.
O Ti (%) = O−0.4Ca−0.66Mg−0.17REM−0.35Zr−0.89Al (5) {However, in the formula (5), component elements that are treated as inevitable impurities are Include in calculation}
Bef (%) = B−0.77 {N−0.29 (Ti−2O Ti )} (6) {However, in the formula (6), when O Ti ≦ 0, O Ti = 0 To do. Further, when O Ti > 0, it is assumed that Ti-2O Ti ≧ 0.005 (%) is satisfied. Further, when N−0.29 (Ti−2O Ti ) ≦ 0 (provided that O Ti ≦ 0 and O Ti = 0), N−0.29 (Ti−2O Ti ) = 0. }
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15 (7) 前記Ｓの含有量が０．０００５〜０．００５％であり、かつ前記Ｏの含有量が０．００１〜０．００４％であり、
さらに、質量％で、Ｃａ：０．０００３〜０．００４％及びＭｇ：０．０００３〜０．００４％のうちの１種又は２種を含有することを特徴とする、請求項６に記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板。The S content is 0.0005 to 0.005%, and the O content is 0.001 to 0.004%,
Furthermore, by mass%, it contains 1 type or 2 types of Ca: 0.0003-0.004% and Mg: 0.0003-0.004%, It is characterized by the above-mentioned. Thick high-strength steel sheet with excellent brittle fracture propagation stopping characteristics and high heat input weld heat-affected zone toughness. さらに、質量％で、Ｎｉ：０．０１〜１％、Ｃｕ：０．０１〜１％、Ｃｒ：０．０１〜１％、Ｍｏ：０．０１〜０．５％、及びＮｂ：０．００３〜０．０３％のうちの１種又は２種以上を含有することを特徴とする、請求項６又は７に記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板。   Furthermore, by mass%, Ni: 0.01-1%, Cu: 0.01-1%, Cr: 0.01-1%, Mo: 0.01-0.5%, and Nb: 0.003 Thickness excellent in brittle fracture propagation stop property and high heat input welding heat affected zone toughness according to claim 6 or 7, characterized by containing one or more of -0.03%. Strength steel plate. さらに、質量％で、ＲＥＭ：０．０００３〜０．０２％及びＺｒ：０．０００３〜０．０２％のうちの１種又は２種を含有することを特徴とする、請求項６〜８のいずれかに記載の脆性破壊伝播停止特性と大入熱溶接熱影響部靭性に優れた厚手高強度鋼板。
Furthermore, by mass%, it contains 1 type or 2 types of REM: 0.0003-0.02% and Zr: 0.0003-0.02% of Claim 6-8 characterized by the above-mentioned. A thick high-strength steel sheet excellent in brittle fracture propagation stopping characteristics and high heat input weld heat-affected zone toughness.
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