JP2005029841A - High strength steel for welded structure having excellent low temperature toughness in high heat input weld part haz, and its production method - Google Patents
High strength steel for welded structure having excellent low temperature toughness in high heat input weld part haz, and its production method Download PDFInfo
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 71
- 239000010959 steel Substances 0.000 title claims abstract description 71
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 17
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical compound [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 claims abstract description 8
- 239000000203 mixture Substances 0.000 claims abstract description 7
- 239000012535 impurity Substances 0.000 claims abstract description 3
- 229910052742 iron Inorganic materials 0.000 claims abstract description 3
- 238000001816 cooling Methods 0.000 claims description 22
- 239000002245 particle Substances 0.000 claims description 22
- 238000003466 welding Methods 0.000 claims description 22
- 238000010438 heat treatment Methods 0.000 claims description 17
- 229910000746 Structural steel Inorganic materials 0.000 claims description 16
- 238000001953 recrystallisation Methods 0.000 claims description 13
- 229910001566 austenite Inorganic materials 0.000 claims description 9
- 230000001186 cumulative effect Effects 0.000 claims description 9
- 230000009467 reduction Effects 0.000 claims description 9
- 238000005098 hot rolling Methods 0.000 claims description 8
- 238000005496 tempering Methods 0.000 claims description 3
- 229910052799 carbon Inorganic materials 0.000 abstract description 4
- 229910052748 manganese Inorganic materials 0.000 abstract description 3
- 229910052760 oxygen Inorganic materials 0.000 abstract description 3
- 229910052751 metal Inorganic materials 0.000 abstract 1
- 239000002184 metal Substances 0.000 abstract 1
- 150000002739 metals Chemical class 0.000 abstract 1
- 229910052717 sulfur Inorganic materials 0.000 abstract 1
- 229910000859 α-Fe Inorganic materials 0.000 description 34
- 239000000463 material Substances 0.000 description 20
- 230000000694 effects Effects 0.000 description 19
- 238000005096 rolling process Methods 0.000 description 15
- 239000010953 base metal Substances 0.000 description 12
- 230000015572 biosynthetic process Effects 0.000 description 9
- 230000007423 decrease Effects 0.000 description 8
- 238000000034 method Methods 0.000 description 8
- 150000003568 thioethers Chemical class 0.000 description 8
- 230000009466 transformation Effects 0.000 description 7
- 238000005516 engineering process Methods 0.000 description 6
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 4
- 230000003749 cleanliness Effects 0.000 description 4
- 150000004767 nitrides Chemical class 0.000 description 4
- 150000001247 metal acetylides Chemical class 0.000 description 3
- 230000002829 reductive effect Effects 0.000 description 3
- 238000003303 reheating Methods 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- 229910052715 tantalum Inorganic materials 0.000 description 3
- 229910052726 zirconium Inorganic materials 0.000 description 3
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 2
- QVGXLLKOCUKJST-UHFFFAOYSA-N atomic oxygen Chemical compound [O] QVGXLLKOCUKJST-UHFFFAOYSA-N 0.000 description 2
- 229910001563 bainite Inorganic materials 0.000 description 2
- 230000015556 catabolic process Effects 0.000 description 2
- 238000004581 coalescence Methods 0.000 description 2
- 238000010276 construction Methods 0.000 description 2
- 238000006731 degradation reaction Methods 0.000 description 2
- 239000006185 dispersion Substances 0.000 description 2
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- 229910000734 martensite Inorganic materials 0.000 description 2
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- 239000001301 oxygen Substances 0.000 description 2
- 239000002244 precipitate Substances 0.000 description 2
- 238000011084 recovery Methods 0.000 description 2
- 238000009628 steelmaking Methods 0.000 description 2
- 238000005728 strengthening Methods 0.000 description 2
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
- 238000004220 aggregation Methods 0.000 description 1
- 230000002776 aggregation Effects 0.000 description 1
- 238000005275 alloying Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 229910052804 chromium Inorganic materials 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 239000002131 composite material Substances 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 239000002826 coolant Substances 0.000 description 1
- 229910052802 copper Inorganic materials 0.000 description 1
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- 239000013078 crystal Substances 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
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- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
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Abstract
Description
【０００１】
【発明の属する技術分野】
本発明は、造船、橋梁、建築、海洋構造物、ラインパイプ、建設機械などの溶接構造物として広く利用可能な、母材靭性と溶接部における熱影響部（以下、ＨＡＺという）の低温靭性の両方に優れた４９０ＭＰａ級の引張強度を有する溶接構造物用鋼およびその製造方法に関するものである。
【０００２】
【従来の技術】
建築、橋梁、造船、海洋構造物など溶接構造物の脆性破壊防止の観点から、母材の靭性だけでなく、溶接部からの脆性破壊の発生抑制すなわち、使用される鋼板のＨＡＺ靱性の向上に関する研究が数多く報告されてきた。一般に、母材靭性の確保のためには最終のフェライト粒径を小さくすることが肝要であり、必要靭性レベルにより普通圧延、制御圧延、さらには制御圧延＋加速冷却が利用されてきた。その基本はＡｌＮやＴｉＮなどの高温で安定な窒化物をピニング粒子として用いて、まず母材の加熱オーステナイト（γ）粒径を微細化した上で、さらに圧延によりオーステナイト中にフェライトの核生成サイトを多数導入し、最終フェライト粒径を微細にすることにある。
【０００３】
したがって、このような母材の製造方法では、当然ながら窒化物の種類により熱間圧延前の再加熱温度を変える必要が生じたり、加熱γ粒径の変動から最終のフェライト粒径にも変化が生じ、結果的に、母材靭性にバラツキが生じることがしばしば起こる。一方、溶接部ＨＡＺ靭性も加熱γ粒径が入熱量によって異なることから、要求靭性値が高いほどその値を小さくする必要があるにも関わらず、近年では加熱γ粒径が大きくなる条件、すなわち溶接施工能率の向上の観点から、大入熱溶接（およそ２０ｋＪ／ｍｍ以下）や超大入熱溶接（２０〜１５０ｋＪ／ｍｍ）が実施される場合が増加している。大入熱溶接と超大入熱溶接の鋼板への影響の差異は、高温での滞留時間の差異に起因しており、特に超大入熱溶接ではその時間が極めて長時間であるために、結晶粒径が著しく粗大化する領域が広く、靱性の低下が著しくなる点にある。
【０００４】
以上のような母材靭性のバラツキと溶接部ＨＡＺ靭性の入熱依存性の問題点を回避する抜本的な方法として、母材組織および溶接部ＨＡＺ組織の加熱γ粒径を同一のピニング粒子によって制御し、両者の高温での粒成長を顕著に抑制することが有効と考えられる。これが実現できた場合は、母材靭性の安定性はもとより入熱が大きくなった場合にも溶接部ＨＡＺ靱性を十分に確保することができる。また、母材の加熱γ粒径が著しく微細になる場合には、従来の制御圧延や加速冷却を用いることなく普通圧延でも同程度のフェライト粒径と母材靭性を付与できる可能性が出てくることから、本技術の確立は工業的価値が高い。
【０００５】
加熱γ粒径のピニング効果が最も期待できる粒子として、高温でも溶解しにくい酸化物や硫化物が考えられる。例えば、酸化物の導入方法としては鋼の溶製工程においてＴｉなどの脱酸元素を単独に添加する方法があるが、多くの場合に溶鋼保持中に酸化物の凝集合体がおこり粗大な酸化物の生成をもたらすことにより、かえって鋼の清浄度を損ない靱性を低下させてしまうことが知られている。そのため、複合脱酸法などさまざまな工夫がなされており、近年ではＴｉ−Ｍｇ脱酸などの方法によって、清浄度を損なうことなく目的とする高温での母材の加熱γ粒径、さらには溶接入熱が大きい場合の加熱γ粒径をかなりの精度で制御できるようになっている。しかしながら、溶接部ＨＡＺの低温靭性については、そのような細粒化効果に加えて、低温においては切り欠き感受性や応力集中度が大きくなるために、いわゆる島状マルテンサイトやセメンタイト（Ｆｅの炭化物）のような硬化相の低減を図る必要があり、低Ｃ化ないしは炭素当量を下げることが不可欠となっている。このようなケースでは、粒界部に優先的に生成される軟質の粒界フェライトの生成挙動が低温靭性に大きく影響を及ぼすことになる。
【０００６】
したがって、ＨＡＺ細粒鋼（例えばＭｇ脱酸等により加熱γ粒径を２００μｍ程度にした場合に相当するが、これはＨＡＺ組織としては極めて細粒であるものの母材に比較するとまだまだ粗大粒に対応する）では、粒界面積が従来鋼に比べて増大しているために、焼入れ性が低下しており、粒界フェライトの生成頻度がより高くなっている。その結果、成分的にも炭素当量が低いことも重なって、フェライト変態が促進され、粗大な粒界フェライトに成り易い。それ故、粒界フェライトの生成抑制あるいは一度核生成した粒界フェライトの成長抑制能力の有無によって靭性ＨＡＺのレベルは大きく変化することになる。実用的には、造船分野などにおいて用いられる大入熱用鋼においてはそのような粒界フェライトの抑制技術が確立されておらず、−４０℃などの過酷な条件において求められる溶接部ＨＡＺの低温靭性を工業的に安定して提供できる鋼板は現状では提案されていない。
【０００７】
【発明が解決しようとする課題】
本発明者らは、酸化物（あるいは硫化物）を最大限に微細分散させた上で、さらに超大入熱溶接時の粒界フェライトの生成抑止技術を鋭意検討し、大入熱や超大入熱溶接においても溶接部ＨＡＺ組織を微細化させ、さらに粒界フェライトの靭性への影響を小さくする技術に着眼し、低温靭性を飛躍的に向上させた高強度溶接構造用鋼の製造技術の確立を課題とした。
【０００８】
【課題を解決するための手段】
上記課題を解決するための本発明の要旨は、以下の通りである。
（１） 質量％で
Ｃ ：０．０３〜０．１４％
Ｓｉ：０．０２〜０．５０％
Ｍｎ：０．３〜２．０％
Ｐ ：≦０．０２％
Ｓ ：０．０００１〜０．０１０％
Ａｌ：０．０００５〜０．０５％
Ｔｉ：０．００３〜０．０５０％
Ｂ ：０．０００１〜０．００２５％
Ｏ ：≦０．０５％
を含み、残部が鉄および不可避的不純物からなり、さらに、質量％で、Ｍｇ：０．０００１〜０．００５０％、Ｃａ：０．０００１〜０．００５０％のうち１種以上を含有し、その合計量が０．００５０％を越えないことを特徴とする大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼。
（２） 質量％で、
Ｃｕ：０．０５〜１．５％
Ｎｉ：０．０５〜５．０％
Ｃｒ：０．０２〜１．５％
Ｍｏ：０．０２〜１．５０％
Ｎｂ：０．０００１〜０．２０％
Ｚｒ：０．０００１〜０．０５０％
Ｔａ：０．０００１〜０．０５０％
のうち１種または２種以上を含有する請求項１記載の大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼。
（３） 溶接部ＨＡＺ組織の加熱γ粒径（旧オーステナイト粒径）が溶接入熱によらず２００μｍ以下であることを特徴とする（１）および（２）記載の大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼。
（４） （１）または（２）記載の鋼と同一成分を有する鋼塊をＡｃ３点以上、１３５０℃以下に加熱後、再結晶温度域で熱間圧延した後、自然冷却することを特徴とする大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼の製造方法。
（５） （１）または（２）記載の鋼と同一成分を有する鋼塊をＡｃ３点以上、１３５０℃以下に加熱後、再結晶温度域で熱間圧延し、さらに未再結晶温度域において累積圧下率で４０〜９０％の熱間圧延をした後、自然冷却することを特徴とする大入熱溶接熱部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼の製造方法。
（６） （１）または（２）記載の鋼と同一成分を有する鋼塊をＡｃ３点以上、１３５０℃以下に加熱後、再結晶温度域で熱間圧延し、さらに未再結晶温度域において累積圧下率で４０〜９０％の熱間圧延をした後、１〜６０℃／ｓｅｃの冷却速度で０〜６００℃まで冷却することを特徴とする大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼の製造方法。
（７） （１）または（２）記載の鋼と同一成分を有する鋼塊をＡｃ３点以上、１３５０℃以下に加熱後、再結晶温度域で熱間圧延し、さらに未再結晶温度域において累積圧下率で４０〜９０％の熱間圧延をした後、１〜６０℃／ｓｅｃの冷却速度で０〜６００℃まで冷却し、引き続いて３００℃〜Ａｃ１点に加熱して焼戻し熱処理することを特徴とする大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼の製造方法。
【０００９】
【発明の実施の形態】
ＭｇとＣａは、従来から強脱酸剤、脱硫剤として鋼の清浄度を高めることで、ＨＡＺ靭性を向上させることが知られている。また、これら元素を含有する酸化物の分散を制御して、母材靭性および溶接部ＨＡＺ靱性の両方を向上させる技術として用いた例が特開２００３−４９２３７号公報に記載されている。
本発明者らは、同じようにＭｇ、Ｃａの強脱酸剤あるいは強力な硫化物生成能に着目し、これら元素の添加順序および量を制御することで、大入熱および超大入熱溶接部のＨＡＺにおける低温靭性を向上させる方法として、加熱γ粒径の微細化だけでなく、脆性破壊の起点として作用する介在物の制御技術としても期待できる余地があり、この技術と従来の粗大な加熱γ粒においてＩＧＦ変態の核としての効力を発揮するＢ析出物（ＢＮあるいはＭ２３（Ｃ，Ｂ）６）を適切に組合せることにより、ＨＡＺ細粒鋼の低温靭性の劣化を誘引する粒界フェライトの伸長化を抑制することが十分に可能と推定し、大入熱溶接時おける溶接部ＨＡＺの低温靭性を顕著に向上させることができるものと考えた。
【００１０】
以下、本発明に関して詳細に説明する。
本発明者らは、Ｔｉを添加し弱脱酸した溶鋼中にＭｇあるいはＣａを添加した場合の酸化物の状態を系統的に調べた。その結果、Ｓｉ、Ｍｎによる脱酸後に、Ｔｉ添加、Ｍｇ（Ｃａ）添加の順に添加した場合に、あるいはＴｉ添加とＭｇ（Ｃａ）添加を同時に行い、さらに平衡状態になった状態で再度Ｍｇ（Ｃａ）を添加するというサイクルを行なうことで、Ｍｇ（あるいはＣａ）の酸化物あるいは硫化物が極めて微細に、かつ高密度に生成されることを見出した。このＭｇ添加の効果はＣａをＭｇの代わりに用いても同様に得られ、いずれの元素を添加した場合も添加元素を含む酸化物もしくは硫化物が生成され、その粒子径は０．００５〜０．５μｍ、粒子数は鋼中に１ｍｍ２当たり１００００個以上であり、強力なピニング力を有していることが確認され、溶接部ＨＡＺ組織の加熱γ粒径が溶接入熱によらず２００μｍ以下となる。
【００１１】
しかも、そのような細粒鋼においてＢを添加した場合には、冷却中に生成される粒界フェライトが明らかにＢフリー鋼に比較して微細化すること、また、入熱量が１０ｋＪ／ｍｍ程度では高温段階でＢが粒界偏析するために粒界での焼入れ性が増大し、粒界フェライトの生成が顕著に抑制されること、さらに入熱量が２０ｋＪ／ｍｍ〜５０ｋＪ／ｍｍの範囲では、通常は粒界に沿って長く伸びた、いわゆる伸長なフェライトになるべきところが、多角形状のポリゴナルフェライトに形態変化し、粒界フェライトの粒子数も格段に多くなる。この機構は、鋼中に存在するＢ析出物［ＢＮやＦｅ２３（Ｃ，Ｂ）６］が溶接時の高温にさらされた状態では一旦溶解するが、その後の冷却段階において加熱γ粒界に優先的に析出し、これを核にして粒界フェライトが生成したことを示すもので、単一の粗大な粒界フェライトに比較して、耐脆性破壊特性が著しく向上することになる。
【００１２】
本発明は上記の介在物の存在状態とＢ添加の効果により達成される母材靭性と溶接部ＨＡＺ靱性の両方に優れた鋼材に関するものであり、特に、ＨＡＺ靭性の場合には−４０℃以下のような厳しい条件下においても十分な低温靭性を付与することができる極めて画期的な技術である。すなわち、本発明の特徴は、母材の加熱γ粒径（旧オーステナイト粒径）が再加熱温度によらず１００μｍ以下であり、同時に溶接部ＨＡＺ組織の加熱γ粒径（旧オーステナイト粒径）も２００μｍ以下の細粒であるのに加えて、低温靭性の劣化の主因である粒界フェライトの生成抑制と、たとえ生成した場合においてもその粒成長を極力抑制することに成功したものであり、これらのミクロ組織によって、母材靭性と溶接部ＨＡＺ靭性の両方に優れた高強度溶接構造用鋼を安価に提供できる点にある。
【００１３】
さらに、本発明では、−２０℃以下の極低温における溶接部ＨＡＺ靭性を確保するためには、ＭｇとＣａの添加量および最終的な鋼中の酸素量（以下Ｏ量と記述）とＳ量の制御が重要な鍵となることも併せて知見した。既に述べたように、ＭｇとＣａの添加については、最初に、Ｓｉ、Ｍｎを添加後、まず、Ｔｉを添加し溶鋼中の酸素量を調整した後、少量のＭｇ（Ｃａでも同様、以下同じ）を徐々に添加するか、あるいは、Ｔｉと少量のＭｇを同時に添加した後に、最終段階で再度Ｍｇを添加する。このような添加順序において、さらにＭｇ添加時のＯ量を制御する意味で、Ａｌのような比較的脱酸力の強い元素による予備脱酸をいずれかのＭｇ添加の前段階で行うことにより、最終のＯ量を正確に制御することが可能となる。特に重要なことは、より過酷な低温域での靭性確保に対しては、脆性破壊の起点として作用する数μｍ以上の酸化物をできうる限り小さくすることが肝要であり、そのためにはＯ量としてピニング効力を発揮するための最小量を確保した上で、数μｍ以上の粒子が増加しないようなレベルにすることが必要条件となる。このＯ量の上限値を種々検討した結果、０．００５％であることが実験的に確認された。本発明では、酸化物と同じように脆性破壊の起点となり得る硫化物に対しても、Ｓ量の上限値として０．０１０％に規定することにより、低温靭性が飛躍的に向上することを見出した。
【００１４】
以下、本発明の成分の限定理由について述べる。
Ｃ：Ｃは鋼における母材強度を向上させる基本的な元素として欠かせない元素であり、その有効な下限として０．０３％以上の添加が必要であるが、０．１４％を越える過剰の添加では、鋼材の溶接性や島状マルテンサイトの生成量を増加させることによりＨＡＺ靱性の低下を招くので、その上限を０．１４％とした。
Ｓｉ：Ｓｉは製鋼上脱酸元素として必要な元素であり、鋼中に０．０２％以上の添加が必要であるが、０．５％を越えるとＨＡＺ靱性を低下させるのでそれを上限とする。
Ｍｎ：Ｍｎは、母材の強度および靱性の確保に必要な元素であるが、２．０％を越えるとＨＡＺ靱性を著しく阻害するが、逆に０．３％未満では、母材の強度確保が困難になるために、その範囲を０．３〜２．０％とする。
【００１５】
Ｐ：Ｐは鋼の靱性に影響を与える元素であり、０．０２％を越えて含有すると鋼材の母材だけでなくＨＡＺの低温靱性を著しく阻害するのでその含有される上限を０．０２％とした。
Ｓ：Ｓは０．０１０％を越えて過剰に添加されると粗大な硫化物の生成の原因となり、低温におけるＨＡＺ靱性を著しく阻害し、一方で、その含有量が０．０００１％未満になると、靭性向上に有効な粒内フェライトの核生成サイトとして作用するＭｎＳ等の硫化物生成量が大幅に少なくなるために、０．０００１〜０．０１０％をその範囲とする。
【００１６】
Ａｌ：Ａｌは通常脱酸剤として添加されるが、本発明においては、Ｍｇ（あるいはＣａ）添加前の予備脱酸元素として用いることにより、ＭｇやＣａの最適な脱酸条件を整えるのに有効な働きをする。しかしながら、Ａｌが０．０５％越えて添加されるとＭｇ、Ｃａの添加効果を阻害するために、これを上限とする。また、Ｍｇ、Ｃａの酸化物を安定に生成するためには０．０００５％は必要であり、これを下限とした。
Ｔｉ：Ｔｉは、脱酸剤として、さらには窒化物形成元素として結晶粒の細粒化に効果を発揮する元素であるが、多量の添加は炭化物の形成による靱性の著しい低下をもたらすために、その上限を０．０５０％にする必要があるが、所定の効果を得るためには０．００３％以上の添加が必要であり、その範囲を０．００３〜０．０５０％とする。
【００１７】
Ｂ：Ｂは一般に、固溶すると焼入れ性を増加させるが、またＢＮとして固溶Ｎを低下させ、溶接部ＨＡＺ靱性を顕著に向上させる元素である。特に、本発明では、粒界フェライトの生成抑制あるいは粒界フェライトの生成を伴う大入熱／超大入熱溶接時には極力その成長を抑えるために利用しており、本発明の重要な元素の一つである。Ｂ量としては、０．０００１％以上の添加でその効果を初めて利用できるようになり、過剰の添加は、靱性の低下を招くために、その上限を０．００２５％とする。
【００１８】
ＯはＭｇやＣａと溶鋼中にて反応し、微細酸化物の生成に不可欠な元素である。最終的な鋼中Ｏ量としては、溶接部ＨＡＺの低温靭性を向上させるためには清浄度の観点からなるべく少なくする方が好ましいが、これはピニング粒子の体積分率を減少させることになるため、両者のバランスの基に制御すべきものである。本発明では、より過酷な温度条件、例えば−４０℃以下のような場合には、Ｏ量が０．００５％を越えると急激に靭性が劣化しはじめる。したがって、Ｏ量の上限値を０．００５％とする。
【００１９】
Ｍｇ：Ｍｇは本発明の主たる合金元素であり、主に脱酸剤あるいは硫化物生成元素として添加されるが、０．００５０％を越えて添加されると、粗大な酸化物あるいは硫化物が生成し易くなり、母材およびＨＡＺ靱性の低下をもたらす。しかしながら、０．０００１％未満の添加では、ピニング粒子として必要な酸化物の生成が十分に期待できなくなるため、その添加範囲を０．０００１〜０．００５０％と限定する。
Ｃａ：Ｃａは硫化物を生成することにより伸長ＭｎＳの生成を抑制し、鋼材の板厚方向の特性、特に耐ラメラティアー性を改善する。さらに、ＣａはＭｇと同様な効果を有していることから、本発明の重要な元素である。Ｃａは０．０００１％未満では、十分な効果が得られないので下限値を０．０００１％にした。逆に、Ｃａが０．００５０％を超えるとＣａの粗大酸化物個数が増加し、超微細な酸化物あるいは硫化物の個数が低下するため、その上限を０．００５０％とする。以上のＭｇとＣａを両方添加する場合には、いずれも強力な脱酸元素であることから、粗大な介在物を生成する危険が大きくなるためその合計量は単独添加の上限値と同レベルの０．００５０％とする。
【００２０】
なお、本発明においては、強度および靱性を改善する元素として、Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏ、Ｎｂ、Ｚｒ、Ｔａの中で、１種または２種以上の元素を必要に応じて添加することができる。
Ｃｕ：Ｃｕは、靱性を低下させずに強度の上昇に有効な元素であるが、０．０５％未満では効果がなく、１．５％を越えると鋼片加熱時や溶接時に割れを生じやすくする。従って、その含有量を０．０５〜１．５％以下とする。
Ｎｉ：Ｎｉは、靱性および強度の改善に有効な元素であり、その効果を得るためには０．０５％以上の添加が必要であるが、５．０％以上の添加では溶接性が低下するために、その上限を５．０％とする。
Ｃｒ：Ｃｒは析出強化による鋼の強度を向上させるために、０．０２％以上の添加が有効であるが、多量に添加すると、焼入れ性を上昇させ、ベイナイト組織を生じさせ、靱性を低下させる。従って、その上限を１．５％とする。
【００２１】
Ｍｏ：Ｍｏは、焼入れ性を向上させると同時に、炭窒化物を形成し強度を改善する元素であり、その効果を得るためには、０．０２％以上の添加が必要になるが、１．５０％を越えた多量の添加は必要以上の強化とともに、靱性の著しい低下をもたらすために、その範囲を０．０２〜０．５０％以下とする。
Ｎｂ：Ｎｂは、炭化物、窒化物を形成し強度の向上に効果がある元素であるが、０．０００１％以下の添加ではその効果がなく、０．２０％を越える添加では、靱性の低下を招くために、その範囲を０．０００１〜０．２０％以下とする。
Ｚｒ、Ｔａ：ＺｒとＴａもＮｂと同様に炭化物、窒化物を形成し強度の向上に効果がある元素であるが、０．０００１％以下の添加ではその効果がなく、０．０５０％を越える添加では、逆に靱性の低下を招くために、その範囲を０．０００１〜０．０５０％以下とする。
【００２２】
上記の成分を含有する鋼は、製鋼工程で溶製後、連続鋳造などを経て再加熱、圧延、冷却処理を施される。この場合、以下の点を限定した。
熱間圧延・制御圧延ともに、鋼塊をオーステナイト化するためにＡｃ３点以上の温度に加熱する必要がある。しかし、１３５０℃を超えて加熱すると、熱源コストの増大が生じることから、加熱温度は１３５０℃以下とした。
次いで、熱間圧延・制御圧延ともに、再結晶温度域で圧延することによりオーステナイト粒径を小さくすることが必要である。また、制御圧延を用いて、強度上昇と靭性向上を図る場合には、さらに未再結晶温度域で圧延することによりオーステナイト粒内に変形帯を導入し、フェライト変態核を導入することが有効である。未再結晶域での累積圧下率が４０％未満では変形帯が十分に形成されないので、未再結晶域で累積圧下率の下限値を４０％とした。しかし、累積圧下率が９０％を超えると、母材シャルピー試験の吸収エネルギーの低下が著しくなるために、上限を９０％にした。
【００２３】
自然放冷よりさらに強度を上昇させるためには加速冷却が必要である。しかしながら、冷却速度が１℃／ｓｅｃ未満では、十分な強度を得ることができない。逆に、冷却速度が６０℃／ｓｅｃ超ではベイナイト主体組織が生成するため母材の靭性が低下する。したがって、冷却速度を１〜６０℃／ｓｅｃに限定した。本発明においては、母材の強度を得るために変態が終了するまで加速冷却を継続する必要がある。このため、冷却停止温度の上限を６００℃とした。６００℃超の停止温度では変態が終了しないために、十分な強度が得られない。通常、加速冷却は水を冷却媒体として用いる。それ故、実際上の冷却停止温度の下限は０℃となるので、下限値を０℃とした。
【００２４】
加速冷却後の焼戻し熱処理は回復による母材組織の靭性向上を目的としたものであるから、加熱温度は逆変態が生じない温度域であるＡｃ１点以下でなければならない。回復は転位の消滅・合体により格子欠陥密度を減少させるものであり、これを実現するためには３００℃以上に加熱することが必要である。このため、加熱温度の下限を３００℃とした。上限は変態点以下であるため、Ａｃ１を上限とした。
【００２５】
【実施例】
表１の化学成分を有する鋼塊を表２に示す製造条件により、板厚１２ｍｍ〜１００ｍｍの厚鋼板とした後、溶接入熱が１０ｋＪ／ｍｍ、２０ｋＪ／ｍｍおよび５０ｋＪ／ｍｍの大入熱溶接あるいは超大入熱溶接を施し、旧γ粒界における粒界フェライトの生成状況をミクロ組織にて調査するとともに、溶接部ＨＡＺ靭性を−４０℃におけるシャルピー吸収エネルギーによって評価した。なお、母材靭性については、加熱温度を１１５０℃と１２５０℃の２水準の温度にて製造しているが、全て良好な母材靭性であった。発明鋼の延性・脆性遷移温度（ｖＴｒｓ）は−４０℃以下であり、試験温度−４０〜−８０℃の範囲にて高いシャルピー吸収エネルギー値（１００Ｊ以上）を示した。
【００２６】
次に、本発明の重要な特性である溶接部ＨＡＺの低温靭性に関しての実施例について説明する。まず、鋼１〜２２は本発明の例を示している。表１、表２から明らかなように、本発明の鋼板は化学成分、製造条件の各要件を満足しており、各入熱における加熱γ粒径はいずれも２００μｍ以下となっていた。さらに、粒界フェライトの生成状況については１０ｋＪ／ｍｍの場合にはほとんど生成されておらず、他方、２０ｋＪ／ｍｍと５０ｋＪ／ｍｍの入熱においてはいずれも粒界フェライトが生成されているにも関わらず、その粒径は最大でも３０μｍであり、かつ伸長の粒界フェライトは皆無であった。
ＨＡＺ靭性はｖＥ−４０，１０ｋとしては２００Ｊ程度の高靭性を示し、また、ｖＥ−４０，２０ｋ、ｖＥ−４０，５０ｋも両方とも１００Ｊを越える値となっており、発明鋼が良好な低温靭性を有していることがわかる。
【００２７】
それに対し、鋼２３〜３５は本発明方法から逸脱した比較例である。すなわち、鋼２３〜２７は基本成分あるいは選択元素の内いずれかの元素が、発明の要件を越えて添加されている例であり、粒界フェライトの生成状況には無関係に靭性が劣位となっている。これは靱性劣化要因となる元素が過剰に添加された事による。また、鋼２８〜３１ではＡｌとＴｉがいずれも下限値ないしは上限値を逸脱した場合に相当している。これらの例では、Ｂ量や脱酸元素であるＭｇとＣａの要件が満足されているものの予備脱酸元素として用いられるＡｌ量とＴｉ量の逸脱も低温靭性に大きく影響を及ぼすことを示しており、これらの条件下ではすべて低靭性となっている。次に、鋼３２〜鋼３４はいずれも重要元素が範囲外の場合であり、鋼３２はＢが高いことに加えて焼入れ性の高いＭｏも高いことから靭性レベルは低い。鋼３３と鋼３４はＭｇ量が０．００５％を越えている場合とＭｇ＋Ｃａの合計量が同じく０．００５％を越えている例である。ちなみに、鋼３３ではＯ量も０．００５０％以上となっている。これらの場合のミクロ組織は一見すると発明鋼と同等の組織、すなわち微細な加熱γ粒径かつ粒界フェライトが著しく微細化されており、高靭性を示しても不思議ではないが、シャルピー試験片の破面観察を行うと脆性破壊の起点として粗大な介在物が多数認められ、そのために靭性が低下したものと理解できる。鋼３５はＢ、Ｍｇ、Ｃａの３元素がいずれも添加されていない場合であり、他に比べて著しく粒界フェライトが粗大化しており、ＨＡＺ靭性が最も悪い。
【００２８】
【表１】
【表２】
【発明の効果】
本発明の化学成分および製造方法に限定し、ＭｇあるいはＣａを適切に添加することで、母材の加熱γ粒径を微細化することができ、さらにＢを有効に利用することと粗大な酸化物や硫化物の生成を抑制することにより、大入熱溶接に関わらずＨＡＺの加熱γ粒径と粒界フェライトの生成抑制と成長抑制を図ることができ、これらの複合効果により母材靭性確保に加えて、大入熱溶接部ＨＡＺの低温靱性を飛躍的に向上させた高強度溶接構造用鋼の製造が可能となる。特に、低温靭性を必要とする造船や橋梁、建築、海洋構造物、ラインパイプ、さらには建設機械などの溶接構造物の脆性破壊に対する安全性が大幅に向上し、産業上の効果は著しく大きい。[0001]
BACKGROUND OF THE INVENTION
The present invention can be widely used as welded structures such as shipbuilding, bridges, buildings, marine structures, line pipes, construction machines, etc., and has low base material toughness and low temperature toughness of a heat affected zone (hereinafter referred to as HAZ) in a welded portion. The present invention relates to a steel for welded structures having a tensile strength of 490 MPa class excellent in both and a method for producing the same.
[0002]
[Prior art]
From the viewpoint of preventing brittle fracture of welded structures such as architecture, bridges, shipbuilding, and offshore structures, not only the toughness of the base metal but also the suppression of the occurrence of brittle fracture from the welded part, that is, the improvement of the HAZ toughness of the steel sheet used Many studies have been reported. In general, it is important to reduce the final ferrite grain size in order to ensure the toughness of the base metal. Depending on the required toughness level, normal rolling, controlled rolling, and controlled rolling + accelerated cooling have been used. The basics are to use high-temperature stable nitrides such as AlN and TiN as pinning particles, first refine the heated austenite (γ) grain size of the base material, and then roll the ferrite nucleation sites in the austenite In order to make the final ferrite grain size fine, a large number of them are introduced.
[0003]
Therefore, in such a manufacturing method of the base material, it is natural that the reheating temperature before hot rolling needs to be changed depending on the type of nitride, or the final ferrite particle size changes due to the variation of the heated γ particle size. Often results in variations in the base material toughness. On the other hand, the weld HAZ toughness also varies depending on the amount of heat input, so that the higher the required toughness value, the smaller the value needs to be reduced in recent years. From the viewpoint of improving the welding work efficiency, cases where large heat input welding (approximately 20 kJ / mm or less) and super large heat input welding (20 to 150 kJ / mm) are increasing. The difference in the effect of high heat input welding and super high heat input welding on the steel sheet is due to the difference in residence time at high temperatures, and especially in super high heat input welding, the time is extremely long. The area where the diameter is remarkably coarsened is wide and the toughness is remarkably lowered.
[0004]
As a fundamental method for avoiding the above-mentioned variation in base metal toughness and heat input dependence of welded part HAZ toughness, the heating γ grain size of the base metal structure and welded part HAZ structure is determined by the same pinning particles. It is considered effective to control and remarkably suppress the grain growth at both high temperatures. If this can be realized, the weld HAZ toughness can be sufficiently secured not only in the stability of the base metal toughness but also in the case where the heat input becomes large. In addition, when the heated γ grain size of the base material becomes extremely fine, there is a possibility that the same ferrite grain size and base material toughness can be imparted even with ordinary rolling without using conventional controlled rolling or accelerated cooling. Therefore, the establishment of this technology has high industrial value.
[0005]
Oxides and sulfides that are difficult to dissolve even at high temperatures can be considered as the particles that are most expected to have the pinning effect of the heated γ particle diameter. For example, as a method for introducing an oxide, there is a method in which a deoxidizing element such as Ti is added alone in the steel melting process. In many cases, however, the oxide is a coarse oxide due to aggregation and coalescence of the oxide during holding of the molten steel. In other words, it is known that the cleanliness of the steel is impaired and the toughness is lowered. For this reason, various devices such as a composite deoxidation method have been made. In recent years, by using a method such as Ti-Mg deoxidation, the heating γ particle size of the base material at a desired high temperature without impairing the cleanliness, and further welding. The heating γ particle size when the heat input is large can be controlled with considerable accuracy. However, with regard to the low temperature toughness of the weld HAZ, in addition to such a fine graining effect, notch sensitivity and stress concentration increase at low temperatures, so-called island martensite and cementite (Fe carbide). Therefore, it is indispensable to reduce the C or lower the carbon equivalent. In such a case, the generation behavior of soft grain boundary ferrite that is preferentially generated at the grain boundary part greatly affects the low temperature toughness.
[0006]
Therefore, HAZ fine-grained steel (corresponding to the case where the heated γ grain size is set to about 200 μm by Mg deoxidation, for example, corresponds to a coarse grain compared with the base material although the HAZ structure is very fine. ), The interfacial area of the grain is increased as compared with that of the conventional steel, so that the hardenability is lowered and the generation frequency of the grain boundary ferrite is higher. As a result, the carbon equivalent is also low, and the ferrite transformation is promoted, so that coarse grain boundary ferrite tends to be formed. Therefore, the level of toughness HAZ varies greatly depending on the presence or absence of the ability to suppress the formation of grain boundary ferrite or the ability to suppress the growth of grain boundary ferrite once nucleated. Practically, in the steel for large heat input used in the shipbuilding field and the like, such a technique for suppressing the grain boundary ferrite has not been established, and the low temperature of the weld HAZ required under severe conditions such as -40 ° C. At present, no steel sheet that can provide toughness stably industrially has been proposed.
[0007]
[Problems to be solved by the invention]
The inventors of the present invention have made fine dispersion of oxides (or sulfides) to the maximum extent, and have further studied the technology for inhibiting the formation of intergranular ferrite during super-high heat input welding. In welding, we will focus on technologies to refine the weld zone HAZ structure and reduce the effect of intergranular ferrite on toughness, and establish high-strength welded structural steel manufacturing technology that dramatically improves low-temperature toughness. It was an issue.
[0008]
[Means for Solving the Problems]
The gist of the present invention for solving the above problems is as follows.
(1) By mass% C: 0.03-0.14%
Si: 0.02 to 0.50%
Mn: 0.3 to 2.0%
P: ≦ 0.02%
S: 0.0001 to 0.010%
Al: 0.0005 to 0.05%
Ti: 0.003 to 0.050%
B: 0.0001 to 0.0025%
O: ≦ 0.05%
The balance is composed of iron and inevitable impurities, and further contains at least one of Mg: 0.0001 to 0.0050% and Ca: 0.0001 to 0.0050% by mass, High strength welded structural steel excellent in low temperature toughness of high heat input weld HAZ, characterized in that the total amount does not exceed 0.0050%.
(2) By mass%
Cu: 0.05 to 1.5%
Ni: 0.05-5.0%
Cr: 0.02 to 1.5%
Mo: 0.02 to 1.50%
Nb: 0.0001 to 0.20%
Zr: 0.0001 to 0.050%
Ta: 0.0001 to 0.050%
The high-strength welded structural steel excellent in low temperature toughness of the high heat input weld HAZ according to claim 1, which contains one or more of them.
(3) The heating γ grain size (former austenite grain size) of the weld zone HAZ structure is 200 μm or less irrespective of the welding heat input, and the high heat input weld zone HAZ according to (1) and (2) High strength welded structural steel with excellent low temperature toughness.
(4) A steel ingot having the same component as the steel described in (1) or (2) is heated to a temperature of Ac 3 or higher and 1350 ° C. or lower, then hot-rolled in a recrystallization temperature range, and then naturally cooled. The manufacturing method of the high strength welded structural steel excellent in the low temperature toughness of the high heat input welded part HAZ.
(5) After heating the steel ingot having the same component as the steel described in (1) or (2) to Ac 3 points or more and 1350 ° C. or less, it is hot-rolled in the recrystallization temperature range, and further in the non-recrystallization temperature range A method for producing high strength welded structural steel excellent in low temperature toughness of a high heat input hot zone HAZ, wherein the steel is naturally cooled after hot rolling at a cumulative rolling reduction of 40 to 90%.
(6) A steel ingot having the same composition as the steel described in (1) or (2) is heated to Ac 3 points or more and 1350 ° C. or less, then hot-rolled in the recrystallization temperature range, and further in the non-recrystallization temperature range After hot rolling with a cumulative rolling reduction of 40 to 90%, cooling to 0 to 600 ° C. at a cooling rate of 1 to 60 ° C./sec is excellent in low temperature toughness of the high heat input weld HAZ Manufacturing method of high strength welded structural steel.
(7) After heating the steel ingot having the same component as the steel described in (1) or (2) to Ac 3 points or more and 1350 ° C. or less, it is hot-rolled in the recrystallization temperature range, and further in the non-recrystallization temperature range After hot rolling with a cumulative rolling reduction of 40 to 90%, it is cooled to 0 to 600 ° C. at a cooling rate of 1 to 60 ° C./sec, and subsequently heated to 300 ° C. to Ac 1 point for tempering heat treatment. The manufacturing method of the steel for high strength welded structures excellent in the low temperature toughness of the high heat input welding part HAZ characterized by these.
[0009]
DETAILED DESCRIPTION OF THE INVENTION
Mg and Ca are conventionally known to improve HAZ toughness by increasing the cleanliness of steel as a strong deoxidizer and desulfurizer. Japanese Patent Laid-Open No. 2003-49237 discloses an example in which the dispersion of oxides containing these elements is controlled to improve both the base metal toughness and the welded portion HAZ toughness.
In the same way, the present inventors pay attention to the strong deoxidizer of Mg and Ca or the strong sulfide generating ability, and control the addition order and amount of these elements, thereby providing high heat input and super high heat input welds. As a method for improving low temperature toughness in HAZ, there is room to expect not only the refinement of the heated γ grain size but also the control technology of inclusions that act as the starting point of brittle fracture. Grain boundaries that induce degradation of low-temperature toughness of HAZ fine-grained steel by appropriately combining B precipitates (BN or M 23 (C, B) 6 ) that exhibit the IGF transformation nucleus in γ grains It was estimated that it was possible to sufficiently suppress the elongation of ferrite, and it was considered that the low temperature toughness of the weld zone HAZ during high heat input welding could be remarkably improved.
[0010]
Hereinafter, the present invention will be described in detail.
The present inventors systematically investigated the state of oxides when Mg or Ca was added to molten steel that was weakly deoxidized by adding Ti. As a result, after deoxidation with Si and Mn, when Ti addition and Mg (Ca) addition are added in this order, or Ti addition and Mg (Ca) addition are performed simultaneously, and Mg (Ca) is again in an equilibrium state. It has been found that by performing a cycle of adding Ca), an oxide or sulfide of Mg (or Ca) is generated extremely finely and with high density. The effect of adding Mg can be obtained in the same manner even when Ca is used instead of Mg. When any element is added, an oxide or sulfide containing the added element is generated, and the particle size is 0.005 to 0. .5 μm, the number of particles is 10,000 or more per 1 mm 2 in the steel, and it is confirmed that the steel has a strong pinning force, and the heated γ particle size of the weld zone HAZ structure is 200 μm or less regardless of welding heat input. It becomes.
[0011]
Moreover, when B is added to such fine-grained steel, the grain boundary ferrite produced during cooling is clearly made finer than that of B-free steel, and the heat input is about 10 kJ / mm. Then, since B segregates at the high temperature stage at the high temperature stage, the hardenability at the grain boundary is increased, the formation of grain boundary ferrite is remarkably suppressed, and the heat input is in the range of 20 kJ / mm to 50 kJ / mm. Normally, what should be a so-called elongated ferrite elongated along the grain boundary is changed to a polygonal polygonal ferrite, and the number of grains of the grain boundary ferrite is remarkably increased. This mechanism is such that B precipitates [BN and Fe 23 (C, B) 6 ] existing in the steel are once dissolved in a state where they are exposed to the high temperature at the time of welding, but in the subsequent cooling stage, they are heated to the heated γ grain boundary. This indicates that the grain boundary ferrite is preferentially precipitated and formed as a nucleus, and the brittle fracture resistance is remarkably improved as compared with a single coarse grain boundary ferrite.
[0012]
The present invention relates to a steel material excellent in both base metal toughness and welded portion HAZ toughness achieved by the presence of the inclusions and the effect of addition of B, and in particular, in the case of HAZ toughness, -40 ° C or lower. This is an extremely innovative technology that can provide sufficient low temperature toughness even under such severe conditions. That is, the feature of the present invention is that the heated γ particle size (old austenite particle size) of the base material is 100 μm or less regardless of the reheating temperature, and at the same time, the heated γ particle size (old austenite particle size) of the weld zone HAZ structure is also In addition to the fine grains of 200 μm or less, they succeeded in suppressing the formation of intergranular ferrite, which is the main cause of the deterioration of low temperature toughness, and suppressing the grain growth as much as possible. With this microstructure, high strength welded structural steel excellent in both base metal toughness and weld zone HAZ toughness can be provided at low cost.
[0013]
Furthermore, in the present invention, in order to ensure the weld zone HAZ toughness at an extremely low temperature of −20 ° C. or less, the addition amount of Mg and Ca, the final oxygen content in the steel (hereinafter referred to as “O content”), and the S content. We also found that control of the system is an important key. As already mentioned, regarding the addition of Mg and Ca, after adding Si and Mn, first, after adding Ti and adjusting the amount of oxygen in the molten steel, a small amount of Mg (same for Ca, and so on) ), Or after adding Ti and a small amount of Mg simultaneously, Mg is added again in the final stage. In such an addition order, in order to further control the amount of O at the time of Mg addition, by performing preliminary deoxidation with an element having a relatively strong deoxidizing power such as Al at the stage before any Mg addition, It becomes possible to accurately control the final amount of O. It is particularly important to minimize the oxide of several μm or more, which acts as a starting point for brittle fracture, in order to ensure toughness in a severer low temperature range. As a necessary condition, it is necessary to secure a minimum amount for exhibiting the pinning effect and to a level that does not increase the number of particles of several μm or more. As a result of various investigations on the upper limit of the amount of O, it was experimentally confirmed to be 0.005%. In the present invention, it is found that the low temperature toughness is drastically improved by defining 0.010% as the upper limit of the amount of S even for the sulfide which can be the starting point of brittle fracture like the oxide. It was.
[0014]
Hereinafter, the reasons for limiting the components of the present invention will be described.
C: C is an indispensable element as a basic element for improving the strength of the base metal in steel. As an effective lower limit, addition of 0.03% or more is necessary, but an excess exceeding 0.14% is required. Addition causes a decrease in the HAZ toughness by increasing the weldability of the steel material and the amount of island martensite produced, so the upper limit was made 0.14%.
Si: Si is an element necessary as a deoxidizing element in steelmaking, and 0.02% or more is necessary to be added to the steel. However, if it exceeds 0.5%, the HAZ toughness is lowered, so that is the upper limit. .
Mn: Mn is an element necessary for securing the strength and toughness of the base material. However, if it exceeds 2.0%, the HAZ toughness is remarkably inhibited, but if it is less than 0.3%, the strength of the base material is secured. Therefore, the range is made 0.3 to 2.0%.
[0015]
P: P is an element that affects the toughness of steel. If it exceeds 0.02%, not only the base material of steel but also the low temperature toughness of HAZ is significantly inhibited, so the upper limit of its content is 0.02% It was.
S: When S is added in excess of 0.010%, it causes coarse sulfides to be formed, which significantly inhibits HAZ toughness at low temperatures, while its content is less than 0.0001%. In order to significantly reduce the amount of sulfides such as MnS that act as nucleation sites of intragranular ferrite effective in improving toughness, the range is 0.0001 to 0.010%.
[0016]
Al: Al is usually added as a deoxidizer, but in the present invention, it is effective for adjusting the optimum deoxidation conditions for Mg and Ca by using it as a preliminary deoxidation element before adding Mg (or Ca). Work nicely. However, if Al is added in an amount exceeding 0.05%, the upper limit is set in order to inhibit the effect of adding Mg and Ca. Further, 0.0005% is necessary to stably produce Mg and Ca oxides, and this is set as the lower limit.
Ti: Ti is an element that exerts an effect on the refinement of crystal grains as a deoxidizer and further as a nitride-forming element. However, a large amount of addition causes a significant decrease in toughness due to the formation of carbides. Although the upper limit needs to be 0.050%, in order to acquire a predetermined effect, addition of 0.003% or more is required, and the range shall be 0.003-0.050%.
[0017]
B: In general, B is an element that increases the hardenability when dissolved, but lowers the solid solution N as BN and significantly improves the weld zone HAZ toughness. In particular, in the present invention, it is used to suppress the growth as much as possible at the time of high heat input / ultra high heat input welding accompanied by generation of grain boundary ferrite or generation of grain boundary ferrite, and is one of the important elements of the present invention. It is. As the amount of B, the effect can be used for the first time by adding 0.0001% or more, and excessive addition causes a decrease in toughness, so the upper limit is made 0.0025%.
[0018]
O reacts with Mg and Ca in molten steel and is an essential element for the production of fine oxides. The final amount of O in steel is preferably as small as possible from the viewpoint of cleanliness in order to improve the low temperature toughness of the weld HAZ, but this will reduce the volume fraction of pinning particles. It should be controlled based on the balance between the two. In the present invention, under more severe temperature conditions, for example, −40 ° C. or less, the toughness starts to deteriorate rapidly when the O content exceeds 0.005%. Therefore, the upper limit value of the O amount is set to 0.005%.
[0019]
Mg: Mg is the main alloying element of the present invention, and is mainly added as a deoxidizer or sulfide-forming element, but if added over 0.0050%, coarse oxides or sulfides are formed. Resulting in a decrease in the base metal and HAZ toughness. However, if the addition is less than 0.0001%, generation of an oxide necessary as pinning particles cannot be sufficiently expected, so the addition range is limited to 0.0001 to 0.0050%.
Ca: Ca suppresses the generation of stretched MnS by generating sulfides, and improves the properties in the thickness direction of the steel material, particularly the lamellar resistance. Furthermore, Ca is an important element of the present invention because it has the same effect as Mg. If Ca is less than 0.0001%, a sufficient effect cannot be obtained, so the lower limit was made 0.0001%. Conversely, when Ca exceeds 0.0050%, the number of coarse oxides of Ca increases and the number of ultrafine oxides or sulfides decreases, so the upper limit is made 0.0050%. When both Mg and Ca are added, since both are strong deoxidizing elements, the risk of generating coarse inclusions increases, so the total amount is at the same level as the upper limit of single addition. 0.0050%.
[0020]
In the present invention, as an element for improving strength and toughness, one or more elements among Cu, Ni, Cr, Mo, Nb, Zr, and Ta may be added as necessary. it can.
Cu: Cu is an element effective in increasing the strength without reducing toughness, but if it is less than 0.05%, it is not effective, and if it exceeds 1.5%, it tends to cause cracking when heating the steel slab or welding. To do. Therefore, the content is made 0.05 to 1.5% or less.
Ni: Ni is an element effective for improving toughness and strength. To obtain the effect, 0.05% or more of addition is necessary. However, addition of 5.0% or more lowers weldability. Therefore, the upper limit is made 5.0%.
Cr: Cr is effective to add 0.02% or more to improve the strength of steel by precipitation strengthening, but if added in a large amount, the hardenability is increased, the bainite structure is generated, and the toughness is reduced. . Therefore, the upper limit is made 1.5%.
[0021]
Mo: Mo is an element that improves hardenability and at the same time forms carbonitride to improve strength. To obtain the effect, addition of 0.02% or more is necessary. The addition of a large amount exceeding 50% brings about a remarkable decrease in toughness as well as an unnecessarily strengthening, so the range is made 0.02 to 0.50% or less.
Nb: Nb is an element that forms carbides and nitrides and is effective in improving the strength. However, the addition of 0.0001% or less has no effect, and the addition exceeding 0.20% reduces toughness. Therefore, the range is made 0.0001 to 0.20% or less.
Zr, Ta: Zr and Ta are elements that form carbides and nitrides as well as Nb and are effective in improving the strength. However, addition of 0.0001% or less has no effect and exceeds 0.050%. Addition, on the contrary, causes a decrease in toughness, so the range is made 0.0001 to 0.050% or less.
[0022]
The steel containing the above components is subjected to reheating, rolling, and cooling through continuous casting after melting in the steelmaking process. In this case, the following points were limited.
In both hot rolling and controlled rolling, it is necessary to heat the steel ingot to a temperature of 3 or more points in order to austenite. However, if heating exceeds 1350 ° C., the heat source cost increases, so the heating temperature is set to 1350 ° C. or lower.
Next, in both hot rolling and controlled rolling, it is necessary to reduce the austenite grain size by rolling in the recrystallization temperature range. In addition, when using controlled rolling to increase strength and improve toughness, it is effective to introduce a deformation band into the austenite grains by rolling in the non-recrystallization temperature range and introduce ferrite transformation nuclei. is there. If the cumulative reduction rate in the non-recrystallized region is less than 40%, the deformation band is not sufficiently formed. Therefore, the lower limit value of the cumulative reduction rate in the non-recrystallized region is set to 40%. However, if the cumulative rolling reduction exceeds 90%, the absorbed energy in the base metal Charpy test is significantly reduced, so the upper limit was made 90%.
[0023]
Accelerated cooling is required to increase the strength further than natural cooling. However, if the cooling rate is less than 1 ° C./sec, sufficient strength cannot be obtained. On the contrary, when the cooling rate exceeds 60 ° C./sec, the toughness of the base material decreases because a bainite main structure is generated. Therefore, the cooling rate was limited to 1-60 ° C./sec. In the present invention, it is necessary to continue accelerated cooling until the transformation is completed in order to obtain the strength of the base material. For this reason, the upper limit of the cooling stop temperature was set to 600 ° C. Since the transformation does not end at a stop temperature exceeding 600 ° C., sufficient strength cannot be obtained. Usually, accelerated cooling uses water as a cooling medium. Therefore, since the lower limit of the actual cooling stop temperature is 0 ° C., the lower limit is set to 0 ° C.
[0024]
Since the tempering heat treatment after accelerated cooling is intended to improve the toughness of the base metal structure by recovery, the heating temperature must be Ac 1 point or less, which is a temperature range in which reverse transformation does not occur. Recovery reduces the lattice defect density by the disappearance and coalescence of dislocations. In order to realize this, heating to 300 ° C. or higher is necessary. For this reason, the minimum of heating temperature was 300 degreeC. Since the upper limit is below the transformation point, Ac 1 was set as the upper limit.
[0025]
【Example】
A steel ingot having the chemical composition shown in Table 1 is made into a thick steel plate having a thickness of 12 mm to 100 mm according to the manufacturing conditions shown in Table 2, and then a large heat input welding with welding heat input of 10 kJ / mm, 20 kJ / mm and 50 kJ / mm. Alternatively, super-high heat input welding was performed, and the formation state of intergranular ferrite at the old γ grain boundary was investigated in the microstructure, and the weld HAZ toughness was evaluated by Charpy absorbed energy at −40 ° C. In addition, about the base material toughness, although the heating temperature was manufactured at the temperature of two levels of 1150 degreeC and 1250 degreeC, all were base material toughness favorable. The ductile / brittle transition temperature (vTrs) of the inventive steel was −40 ° C. or lower, and showed a high Charpy absorbed energy value (100 J or higher) in the range of the test temperature from −40 to −80 ° C.
[0026]
Next, examples relating to the low temperature toughness of the weld HAZ, which is an important characteristic of the present invention, will be described. First, steels 1-22 show examples of the present invention. As is apparent from Tables 1 and 2, the steel sheet of the present invention satisfied the requirements of chemical components and production conditions, and the heating γ particle size at each heat input was 200 μm or less. Further, regarding the generation state of the grain boundary ferrite, it is hardly generated in the case of 10 kJ / mm, and on the other hand, in both heat input of 20 kJ / mm and 50 kJ / mm, the grain boundary ferrite is generated. Regardless, the maximum grain size was 30 μm and there was no elongation grain boundary ferrite.
HAZ toughness as a vE -40,10K exhibit high toughness of about 200 J, also, vE -40,20k, vE -40,50k also has a both value exceeding 100 J, invention steel good low-temperature toughness It can be seen that
[0027]
On the other hand, steels 23 to 35 are comparative examples deviating from the method of the present invention. That is, steels 23 to 27 are examples in which any one of the basic components or selected elements is added beyond the requirements of the invention, and the toughness is inferior regardless of the state of formation of grain boundary ferrite. Yes. This is due to the excessive addition of elements that cause toughness degradation. Moreover, in steel 28-31, it corresponds when both Al and Ti deviate from the lower limit value or the upper limit value. In these examples, although the requirements for the B content and the deoxidation elements Mg and Ca are satisfied, the deviation of the Al content and the Ti content used as the preliminary deoxidation elements also greatly affects the low temperature toughness. Under these conditions, all have low toughness. Next, steel 32 to steel 34 are all cases where the important elements are out of the range, and steel 32 has a low toughness because it has high B and also has high hardenability Mo. Steel 33 and steel 34 are examples in which the Mg amount exceeds 0.005% and the total amount of Mg + Ca similarly exceeds 0.005%. Incidentally, the amount of O in the steel 33 is 0.0050% or more. At first glance, the microstructure in these cases is equivalent to that of the invented steel, that is, the fine heated γ grain size and the grain boundary ferrite are remarkably refined, and it is not surprising that it exhibits high toughness. When the fracture surface is observed, it can be understood that a large number of coarse inclusions are observed as the starting point of brittle fracture, and that the toughness is lowered. Steel 35 is a case where none of the three elements B, Mg, and Ca is added. Grain boundary ferrite is remarkably coarsened compared to the others, and the HAZ toughness is the worst.
[0028]
[Table 1]
[Table 2]
【The invention's effect】
By limiting the chemical components and the production method of the present invention and adding Mg or Ca appropriately, the heated γ particle size of the base material can be refined, and B can be used effectively and coarsely oxidized. Suppressing the formation of oxides and sulfides can suppress the formation of HAZ heated γ grain size and intergranular ferrite and suppress growth regardless of high heat input welding. In addition to this, it is possible to produce high strength welded structural steel with dramatically improved low temperature toughness of the high heat input weld HAZ. In particular, the safety against brittle fracture of welded structures such as shipbuilding and bridges, buildings, marine structures, line pipes, and construction machines that require low temperature toughness is greatly improved, and the industrial effect is remarkably great.
Claims (7)
Ｃ ：０．０３〜０．１４％
Ｓｉ：０．０２〜０．５０％
Ｍｎ：０．３〜２．０％
Ｐ ：≦０．０２％
Ｓ ：０．０００１〜０．０１０％
Ａｌ：０．０００５〜０．０５％
Ｔｉ：０．００３〜０．０５０％
Ｂ ：０．０００１〜０．００２５％
Ｏ ：≦０．０５％
を含み、残部が鉄および不可避的不純物からなり、さらに、質量％で、Ｍｇ：０．０００１〜０．００５０％、Ｃａ：０．０００１〜０．００５０％のうち１種以上を含有し、その合計量が０．００５０％を越えないことを特徴とする大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼。C: 0.03 to 0.14% by mass%
Si: 0.02 to 0.50%
Mn: 0.3 to 2.0%
P: ≦ 0.02%
S: 0.0001 to 0.010%
Al: 0.0005 to 0.05%
Ti: 0.003 to 0.050%
B: 0.0001 to 0.0025%
O: ≦ 0.05%
The balance is composed of iron and inevitable impurities, and further contains at least one of Mg: 0.0001 to 0.0050% and Ca: 0.0001 to 0.0050% by mass, High strength welded structural steel excellent in low temperature toughness of high heat input weld HAZ, characterized in that the total amount does not exceed 0.0050%.
Ｃｕ：０．０５〜１．５％
Ｎｉ：０．０５〜５．０％
Ｃｒ：０．０２〜１．５％
Ｍｏ：０．０２〜１．５０％
Ｎｂ：０．０００１〜０．２０％
Ｚｒ：０．０００１〜０．０５０％
Ｔａ：０．０００１〜０．０５０％
のうち１種または２種以上を含有する請求項１記載の大入熱溶接部ＨＡＺの低温靭性に優れた高強度溶接構造用鋼。% By mass
Cu: 0.05 to 1.5%
Ni: 0.05-5.0%
Cr: 0.02 to 1.5%
Mo: 0.02 to 1.50%
Nb: 0.0001 to 0.20%
Zr: 0.0001 to 0.050%
Ta: 0.0001 to 0.050%
The high-strength welded structural steel excellent in low temperature toughness of the high heat input weld HAZ according to claim 1, which contains one or more of them.
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| JP2006241510A (en) * | 2005-03-02 | 2006-09-14 | Nippon Steel Corp | Steel for high strength welded structure having excellent low temperature toughness in high heat input weld haz and its production method |
| JP2016008341A (en) * | 2014-06-25 | 2016-01-18 | 新日鐵住金株式会社 | High-tensile strength steel for welding |
| CN108715976A (en) * | 2018-05-25 | 2018-10-30 | 山东钢铁股份有限公司 | A kind of Ti-Zr-C granules reinforced wear resistances steel and preparation method thereof |
| CN111936654A (en) * | 2018-03-30 | 2020-11-13 | 日铁不锈钢株式会社 | Ferritic stainless steel having excellent ridging resistance |
| WO2023284128A1 (en) * | 2021-07-12 | 2023-01-19 | 南京钢铁股份有限公司 | Low-cost high energy welding-resistant 420 mpa-grade bridge steel and production method therefor |
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2003
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Cited By (6)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2006241510A (en) * | 2005-03-02 | 2006-09-14 | Nippon Steel Corp | Steel for high strength welded structure having excellent low temperature toughness in high heat input weld haz and its production method |
| JP2016008341A (en) * | 2014-06-25 | 2016-01-18 | 新日鐵住金株式会社 | High-tensile strength steel for welding |
| CN111936654A (en) * | 2018-03-30 | 2020-11-13 | 日铁不锈钢株式会社 | Ferritic stainless steel having excellent ridging resistance |
| CN111936654B (en) * | 2018-03-30 | 2022-01-18 | 日铁不锈钢株式会社 | Ferritic stainless steel having excellent ridging resistance |
| CN108715976A (en) * | 2018-05-25 | 2018-10-30 | 山东钢铁股份有限公司 | A kind of Ti-Zr-C granules reinforced wear resistances steel and preparation method thereof |
| WO2023284128A1 (en) * | 2021-07-12 | 2023-01-19 | 南京钢铁股份有限公司 | Low-cost high energy welding-resistant 420 mpa-grade bridge steel and production method therefor |
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