JP4348102B2 - 490 MPa class high strength steel excellent in high temperature strength and manufacturing method thereof - Google Patents
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Description
【０００１】
【発明の属する技術分野】
本発明は、建築、土木、海洋構造物、造船、貯槽タンクなどの一般的な構造物に用いる６００℃以上８００℃以下の温度範囲において、１時間程度の比較的短時間における高温強度が優れた低合金炭素添加の建築構造用高張力鋼（鋼板、鋼管、形鋼、線材）の製造方法に関する。
【０００２】
【従来の技術】
例えば、建築、土木などの分野においては、各種建築用鋼材として、ＪＩＳ等で規格化された鋼材等が広く利用されている。なお、一般の建築構造用鋼材は、約３５０℃から強度低下するため、その許容温度は約５００℃となっている。
【０００３】
すなわち、ビルや事務所、住居、立体駐車場などの建築物に前記の鋼材を用いた場合は、火災における安全性を確保するため、十分な耐火被覆を施すことが義務付けられており、建築関連諸法令では、火災時に鋼材温度が３５０℃以上にならないように規定されている。
【０００４】
これは、前記鋼材では、３５０℃程度で耐力が常温の２／３程度になり、必要な強度を下回るためである。鋼材を建造物に利用する場合、火災時において鋼材の温度が３５０℃に達しないように耐火被覆を施して使用される。そのため、鋼材費用に対して耐火被覆工費が高額となり、建設コストが大幅に上昇することが避けられない。
【０００５】
上記の課題を解決するため、高温耐力を備えた耐火鋼が開発されている。
【０００６】
６００℃以上での高温強度がある鋼の場合、一般に耐火鋼と呼称されており、６００℃で常温降伏強度の２／３以上の高温強度を有する耐火鋼（例えば、特許文献１参照）や、７００℃で高温強度が優れた耐火鋼（例えば、特許文献２参照）が提案されている。その他の６００℃耐火鋼に関する発明の例でも、６００℃での降伏強度を常温降伏強度の２／３以上とすることが一般的となっている。
【０００７】
しかしながら、７００℃耐火鋼、８００℃耐火鋼は、現時点では高温強度の設定（常温降伏強度との比率）に一般則が見られない。例えば、特許文献１では、相当量のＭｏとＮｂを添加した鋼で、６００℃の耐力が常温耐力の７０％以上を確保するものであるが、７００℃〜８００℃の耐力は示されていない。また、６００℃の耐力が常温耐力の７０％程度では、火災時の温度上昇を考慮すると、耐火被覆量の低減は可能であるものの、省略が可能となる建造物は立体駐車場やアトリウムなどの開放的空間に限定されるため、無耐火被覆での使用は著しく限定される。
【０００８】
また、特許文献２では、相当量のＭｏとＮｂを添加した鋼でミクロ組織をベイナイトとすることにより、７００℃の耐力が常温耐力の５６％以上を確保するものであるが、８００℃の耐力は示されていない。
【０００９】
すなわち、これらの例のように６００℃程度の高温強度を確保した鋼は、すでに市場でも使用されており、７００℃で一定の強度を確保する鋼材の発明がなされているが、７００℃〜８００℃での高温強度を確保できる実用鋼の安定的な製造は困難であった。
【００１０】
【特許文献１】
特開平２−７７５２３号公報
【特許文献２】
特開平１０−６８０４４号公報
【００１１】
【発明が解決しようとする課題】
前述のように建築物に鋼材を利用する場合、通常の鋼では高温強度が低いため、無被覆や耐火被覆軽減で利用することができず、高価な耐火被覆を施さなければならなかった。
【００１２】
また、新しく開発された鋼でも、耐火温度は６００〜７００℃までの保証が限界であり、７００℃〜８００℃での無耐火被覆使用およびこれによる耐火被覆工程の省略が可能となる鋼材の開発が望まれていた。
【００１３】
本発明の目的は６００℃以上８００℃以下の温度範囲における高温強度及び溶接性に優れた高張力鋼及び当該鋼を工業的に安定して供給可能な製造方法を提供することにある。
【００１４】
【課題を解決するための手段】
本発明は前述の課題を克服するために、ミクロ組織と添加合金元素量等を最適範囲とすることで目的を達成したもので、その要旨は以下に示す通りである。
【００１５】
（１） 鋼成分が質量％で、
Ｃ：０．００５％以上０．０４％未満、
Ｓｉ：０．５％以下、
Ｍｎ：０．１％以上０．５％以下、
Ｐ：０．０２％以下、
Ｓ：０．０１％以下、
Ｍｏ：０．３〜１．５％、
Ｎｂ：０．０３〜０．１５％、
Ｂ：０．０００５〜０．００３％、
Ａｌ：０．０６％以下、
Ｎ：０．００６％以下、
かつ、残部が鉄および不可避的不純物からなることを特徴とする高温強度に優れた４９０ＭＰａ級高張力鋼。
【００１６】
（２） 質量％で、更に、
Ｎｉ：０．０５〜１．０％、
Ｃｕ：０．０５〜１．０％、
Ｃｒ：０．０５〜１．０％、
Ｔｉ：０．００５〜０．０２５％、
Ｖ：０．０１〜０．１％
の範囲で１種または２種以上を含有することを特徴とする上記（１）に記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００１７】
（３） 質量％で、更に、
Ｃａ：０．０００５〜０．００４％、
ＲＥＭ：０．０００５〜０．００４％、
Ｍｇ：０．０００１〜０．００６％
のいずれか１種または２種以上を含有することを特徴とする上記（１）または（２）記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００１８】
（４） 常温の降伏応力により高温時の降伏応力を無次元化した高温常温降伏応力比ｐ（＝高温降伏応力／常温降伏応力）が、鋼材温度Ｔ（℃）が６００℃以上８００℃以下の範囲で、ｐ≧−０．００３３×Ｔ＋２．８０を満足することを特徴とする上記（１）〜（３）の内のいずれかに記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００１９】
（５） 常温においてフェライト及びベイナイトの混合組織であり、火災相当の高温加熱時に、オーステナイトに逆変態する温度（Ａｃ１）が８００℃超であることを特徴とする上記（１）〜（４）の内のいずれかに記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００２０】
（６） フェライト及びベイナイトの混合母相組織中で高温において熱力学的に安定な炭窒化析出相をモル分率にて５×１０- ４以上保持するとともに、ＢＣＣ相中に固溶するＭｏ、Ｎｂの合計量がモル濃度にて２×１０- ３以上であることを特徴とする上記（５）に記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００２１】
（７） フェライト及びベイナイトの混合母相組織中で高温において熱力学的に安定な炭窒化析出相をモル分率にて５×１０-4以上保持するとともに、ＢＣＣ相中に固溶するＭｏ、Ｎｂ、Ｖ、Ｔｉの合計量がモル濃度にて２×１０-3以上であることを特徴とする上記（５）に記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００２２】
（８） フェライトとベイナイトの混合組織として、ベイナイトの分率が２０〜９０％であることを特徴とする上記（１）〜（７）のいずれかに記載の高温強度に優れた４９０ＭＰａ級高張力鋼。
【００２３】
（９） 上記（１）〜（３）の内のいずれか１項に記載の鋼成分からなり、
ＰCM＝Ｃ＋Ｓｉ／３０＋Ｍｎ／２０＋Ｃｕ／２０＋Ｎｉ／６０＋Ｃｒ／２０＋
Ｍｏ／１５＋Ｖ／１０＋５Ｂ
と定義する溶接割れ感受性組成ＰCMが０．１８％以下であることを特徴とする高温強度及び溶接性に優れた４９０ＭＰａ級高張力鋼。
【００２４】
（１０） 上記（４）〜（７）の内のいずれか１項に記載の特徴を有するとともに、フェライトとベイナイトの混合組織として、ベイナイトの分率が２０〜９０％であることを特徴とする上記（９）記載の高温強度及び溶接性に優れた４９０ＭＰａ級高張力鋼。
【００２５】
（１１） 旧オーステナイト粒の平均円相当径が１２０μｍ以下であることを特徴とする上記（９）または（１０）記載の高温強度及び溶接性に優れた４９０ＭＰａ級高張力鋼。
【００２６】
（１２） 上記（１）〜（３）のいずれか１項に記載の鋼成分からなる鋼片または鋳片を１１００〜１２５０℃の温度範囲に再加熱後、１１００℃以下での累積圧下量を３０％以上として、８５０℃以上の温度で圧延し、ミクロ組織をフェライトとベイナイトの混合組織とすることを特徴とする高温強度に優れた４９０ＭＰａ級高張力鋼の製造方法。
【００２７】
（１３） 上記（１）〜（３）のいずれか１項に記載の鋼成分からなる鋼片または鋳片を１１００〜１２５０℃の温度範囲に再加熱後、１１００℃以下での累積圧下量を３０％以上として、８５０℃以上の温度で圧延し、圧延終了後８００℃以上の温度から６５０℃以下の温度までの冷却速度を０．３Ｋｓ-1以上として、ミクロ組織をフェライトとベイナイトの混合組織とすることを特徴とする上記（４）〜（７）のいずれか１項に記載の高温強度に優れた４９０ＭＰａ級高張力鋼の製造方法。
【００２８】
（１４） 上記（１）〜（３）のいずれか１項に記載の鋼成分からなる鋳片を１１００〜１２５０℃の温度範囲に再加熱後、１１００℃以下での累積圧下量を３０％以上として、８５０℃以上の温度で圧延し、ミクロ組織をフェライトとベイナイトの混合組織とすることを特徴とする高温強度及び溶接性に優れた４９０ＭＰａ級高張力鋼の製造方法。
【００２９】
（１５） 上記（１）〜（３）のいずれか１項に記載の鋼成分からなる鋳片を１１００〜１２５０℃の温度範囲に再加熱後、１１００℃以下での累積圧下量を３０％以上として、８５０℃以上の温度で圧延し、圧延終了後８００℃以上の温度から６５０℃以下の温度までの冷却速度を０．３Ｋｓ−１以上として、ミクロ組織をフェライトとベイナイトの混合組織とすることを特徴とする高温強度及び溶接性に優れた低降伏比４９０ＭＰａ級高張力鋼の製造方法。
【００３０】
【発明の実施の形態】
以下、本発明の詳細を説明する。
【００３１】
本発明者らはすでに、６００℃、７００℃の高温強度が優れた鋼を見出した。６００℃の高温強度が優れた鋼はすでに建築分野で使用されているが、市場では更に高温に耐える鋼への極めて強い要求がある。
【００３２】
高温強度増加に対しては、Ｍｏ、Ｎｂの複合添加により高温にて安定な炭窒化物の析出を促進するとともに、ミクロ組織のベイナイト化により転位密度の増大し、さらには固溶Ｍｏ及びＮｂにより転位回復の遅延を図ることが有効である。しかし、硬質ベイナイトの分率が過剰であると、常温の強度が過大となるため、所要の常温強度に応じて、ミクロ組織を適切なベイナイト分率を有するフェライトとベイナイトの混合組織とする。適切なミクロ組織を造り込み、所要の常温強度範囲を達成するには低Ｃ化が有効である。低Ｃ化は、フェライトとベイナイトの混合母相組織の高温における熱力学的安定性を高め、オーステナイトへの逆変態温度（Ａｃ１）を上昇させる効果も持つ。しかし、この場合、ミクロ組織及び材質が圧延条件とその後の冷却条件により影響を受けやすく、安定的な製造が困難であることが判明した。ミクロ組織制御と高温強度の増加に取り組んだ結果、適量のＢ添加が製造安定化に有効であることを知見し、本発明に至った。
【００３３】
一般的な溶接構造用鋼として、溶接性や低降伏強度比等の特性は、従来と同様に具備する必要があるため、７００℃〜８００℃の高温強度が優れた鋼は極めて困難な課題であった。この課題を解決するため、本発明者らは鋭意検討し、７００℃、８００℃の高温強度はＭｏ、Ｎｂ、Ｖ、Ｔｉ等の合金元素の複合添加による析出強化とミクロ組織のベイナイト化による転位密度の増大、さらには固溶Ｍｏ、Ｎｂ、Ｖ、Ｔｉによる転位回復遅延が有効であり、Ｔｉも若干の効果があることを突き止めた。従って、７００℃〜８００℃の強度と常温の強度、常温と高温の強度比（ＹＳ比＝高温強度／常温強度）の全てを同時に確保するためには、ミクロ組織を適切なフェライトとベイナイトの混合組織とするとともに、添加合金元素量を最適範囲として、高温における母相組織の熱的安定性と適切な整合析出強化効果及び転位回復遅延効果を得ることが重要であることを見出した。
【００３４】
鋼材の降伏強度は、一般に４５０℃近傍から急激に低下するが、これは、温度上昇に伴い熱活性化エネルギーが低下し、転位のすべり運動に対して低温では有効であった抵抗が無効となるためである。本発明者らはＭｏ、Ｎｂ、Ｖ、Ｔｉの複合炭窒化物は、転位のすべり運動に対して６００℃程度の高温まで有効な抵抗として作用することを見出した。さらに、ＢＣＣ相中に固溶したＭｏ、Ｎｂ、Ｖ、Ｔｉは、転位回復遅延に対して有効であり、降伏強度の急激な低下が始まる温度を高温化する効果を持つことを知見するに至った。したがって、７００℃〜８００℃において、鋼材温度をＴ（℃）として、高温常温降伏応力比ｐ（＝高温降伏応力／常温降伏応力）が、ｐ≧−０．００３３×Ｔ＋２．８０を満足する、すなわち、降伏応力比がそれぞれ４９％、１６％以上となるためには、当該温度における鋼中に成分元素として含有されるＭｏ、Ｎｂ、Ｖ、Ｔｉの複合炭窒化物はモル分率にて５×１０-4以上であるとともに、ＢＣＣ相中に固溶する成分元素であるＭｏ、Ｎｂ、Ｖ、Ｔｉの合計量がモル濃度にて２×１０-3以上でなければならない。
【００３５】
高温強度発現に重要である複合炭窒化析出相の組成は、例えば電子顕微鏡やＥＤＸによる分析により容易に同定可能である。
【００３６】
また、熱力学的に安定な析出相の平衡生成量及びＢＣＣ相中の固溶合金元素量については、市販の熱力学計算データベースソフト等利用することにより、添加合金元素量より容易に算出可能である。
【００３７】
ミクロ組織におけるフェライトの分率が増加し、ベイナイトの分率が２０％未満に低下すると、常温及び高温の強度が低下し、Ｍｏ、Ｎｂ、Ｔｉ、Ｖ等の合金元素をより多く添加する必要が生じる。ミクロ組織に占めるフェライトの分率が過大となると、添加合金元素の増加による常温及び高温の強度確保は困難になる。逆にミクロ組織におけるフェライトの分率が低下し、ベイナイトの分率が増加すると常温及び高温の強度が上昇し、Ｍｏ、Ｎｂ、Ｔｉ、Ｖ等の合金元素添加量を低減する必要が生じる。さらに、ミクロ組織に占めるベイナイト分率が９０％を超え過大となると、高温強度については達成可能であるが、常温強度の上昇、ＨＡＺ靭性の劣化、溶接性の劣化が顕著となり添加合金元素の低減による所要の強度範囲、ＨＡＺ靭性及び溶接性の確保が困難となる。
【００３８】
このため、本発明鋼ではミクロ組織をフェライトとベイナイトの混合組織とし、ベイナイトの分率を２０％〜９０％の範囲内とするが、３０〜７０％の範囲内とすることが好ましい。
【００３９】
本発明者らは、ミクロ組織をフェライトとベイナイトの混合組織とし、かつ、ベイナイト分率を安定的に２０〜９０％の範囲に保つ方法について検討し、適量のＢ添加が必須であることを見出した。
【００４０】
本発明が、請求項の通りに鋼成分および製造方法を限定した理由について説明する。
【００４１】
常温と高温の強度を同時に確保するためには、相当量の合金元素の添加が必要であり、４９０ＭＰａ級高張力鋼では、Ｍｏ：０．３〜１．５％、Ｎｂ：０．０３〜０．１５％が必要である。
【００４２】
また、高温強度の向上に対して、さらに、Ｔｉ：０．００５〜０．０２５％、Ｖ：０．０１〜０．１％の範囲の添加が有効である。
【００４３】
Ｍｏ、Ｎｂ、Ｔｉ、V等は主に高温強度の確保のためであり、ＳｉとＭｎの範囲限定は常温強度を所定の範囲に抑制するためである。
【００４４】
鋼の加熱温度はＭｏ、Ｎｂ、Ｔｉ、Ｖをできるだけ固溶状態とするために高い温度が望ましいが、母材の靭性確保の観点から１１００〜１２５０℃に限定した。
【００４５】
圧延終了温度を１１００℃以下で８５０℃以上としたのは、低温域の圧下でＭｏ、Ｎｂ、Ｔｉ、Ｖが炭化物として析出するため８５０℃が下限の温度であり、また、１１００℃を超える温度で圧延を終了すると靭性が不足するためである。
【００４６】
なお、本発明鋼を製造後、脱水素などの目的でＡｃ１変態点以下の温度に再加熱しても、本発明鋼の特徴は何ら損なわれることはない。
【００４７】
次に、本説明に関わるその他の成分元素とその添加量について説明する。
【００４８】
Ｃは、鋼材の特性に最も顕著な効果を及ぼすもので、狭い範囲に制御されなければならない。０．００５以上０．０４％未満が限定範囲である。０．００５％未満のＣ量では強度が不足し、０．０４％以上となると圧延終了後の冷却速度が過大の場合はベイナイトの生成分率が増加し強度が超過、逆に冷却速度が過小の場合はベイナイトの生成分率が低下して強度が不足する。さらに、火災相当の高温加熱時に、フェライトとベイナイトの混合母相組織を熱力学的に安定に保ち、Ｍｏ、Ｎｂ、Ｖ、Ｔｉの複合炭窒化析出物との整合性を維持して、強化効果を確保する上でもＣを０．０４％未満とする必要がある。
【００４９】
Ｓｉは、脱酸上鋼に含まれる元素であり、置換型の固溶強化作用を持つことから常温での母材強度向上に有効であるが、特に６００℃超の高温強度を改善する効果はない。また、多く添加すると溶接性、ＨＡＺ靭性が劣化するため、上限を０．５％に限定した。鋼の脱酸はＴｉ、Ａｌのみでも可能であり、ＨＡＺ靭性、焼入性などの観点から低いほど好ましく、必ずしも添加する必要はない。
【００５０】
Ｍｎは、強度、靭性を確保する上で不可欠な元素ではあるが、置換型の固溶強化元素であるＭｎは、常温での強度上昇には有効であるが、特に６００℃超の高温強度にはあまり大きな改善効果はない。したがって、本発明のような比較的多量のＭｏを含有する鋼において溶接性向上すなわちＰCM低減の観点から０．９％以下とする必要がある。さらに、建築用４９０ＭＰａ級高張力鋼に対しては、常温強度上限を考慮し、０．５％以下に限定した。Ｍｎの上限を低く抑えることにより、連続鋳造スラブの中心偏析の点からも有利となる。なお、下限については、母材の強度、靭性調整上、０．１％以上の添加が必要である。したがって、Ｍｎは０．１〜０．５％の範囲とした。
【００５１】
また、常温の降伏強度及び引張り強度を４９０ＭＰａ級高張力鋼の所要範囲とするためには、圧延終了後８００℃以上の温度から６５０℃以下の温度までの冷却速度を０．３Ｋｓ-1以上とする必要がある。すなわち、約２５ｍｍ未満の比較的薄い鋼板は空冷または加速冷却（水冷）プロセスにて、約２５ｍｍ超の比較的厚い鋼板は加速冷却（水冷）プロセスを適用して製造することができる。
【００５２】
Ｐは、本発明鋼においては不純物であり、Ｐ量の低減はＨＡＺにおける粒界破壊を減少させる傾向があるため、少ないほど好ましい。含有量が多いと母材、溶接部の低温靭性を劣化させるため上限を０．０２％とした。
【００５３】
Ｓは、Ｐと同様本発明鋼においては不純物であり、母材の低温靭性の観点からは少ないほど好ましい。含有量が多いと母材、溶接部の低温靭性を劣化させるため上限を０．０１％とした。
【００５４】
Ｍｏは、７００℃、８００℃の高温強度を確保する上で必要不可欠の元素で、本発明においては最も重要な元素の一つである。高温強度のみの考慮であれば、下限の緩和は可能であるが、後述する低降伏比化の観点から、フェライト＋オーステナイトの二相域熱処理、およびその後必要に応じ焼き戻しを行っても、なお常温での高強度、高靭性を確保するため、下限を０．３％とした。一方、１．５％超の添加は、母材材質の制御（ばらつきの制御や靭性の劣化）が困難になるとともに、経済性を失するため、０．３〜１．５％が限定範囲である。
【００５５】
Ｎｂは、Ｍｏを比較的多量添加する本発明においては、７００℃、８００℃の高温強度を確保するために重要な役割を演ずる元素である。まず、一般的な効果として、オーステナイトの再結晶温度を上昇させ、熱間圧延時の制御圧延の効果を最大限に発揮する上で有用な元素である。また、圧延に先立つ再加熱や焼きならしや焼き入れ時の加熱オーステナイトの細粒化にも寄与する。さらに、析出強化及び転位回復抑制による高温強度向上効果を有し、Ｍｏとの複合添加により高温強度向上に寄与する。０．０３％未満では７００℃及び８００℃における析出硬化及び転位回復抑制の効果が少なく、０．１５％を超えると添加量に対し効果の度合いが減少し、経済的にも好ましくない。また、溶接時の靭性も低下する。よって０．０３〜０．１５％が限定範囲である。
【００５６】
Ｂは、ベイナイトの生成分率を介して強度を制御する上で極めて重要である。すなわち、Ｂはオーステナイト粒界に偏析してフェライトの生成を抑制することを介して焼入性を向上させ、空冷のような冷却速度が比較的小さい場合においてもベイナイトを安定的に生成させるのに有効である。この効果を享受するため、最低０．０００５％以上必要である。しかし、多すぎる添加は焼入性向上効果が飽和するだけでなく、旧オーステナイト粒界の脆化や靭性上有害となるＢ析出物を形成する可能性があるため、上限を０．００３％とした。なお、タンク用鋼などとして、応力腐食割れが懸念されるケースでは、母材および溶接熱影響部の硬さの低減がポイントとなることが多く（例えば、硫化物応力腐食割れ（ＳＣＣ）防止のためにはＨＲＣ≦２２（ＨＶ≦２４８）が必須とされる）、そのようなケースでは焼入性を増大させる過剰なＢ添加は好ましくない。
【００５７】
Ａｌは、一般に脱酸上鋼に含まれる元素であるが、脱酸はＳｉまたはＴｉだけでも十分であり、本発明鋼においては、その下限は限定しない（０％を含む）。しかし、Ａｌ量が多くなると鋼の清浄度が悪くなるだけでなく、溶接金属の靭性が劣化するので上限を０．０６％とした。
【００５８】
Ｎは、不可避的不純物として鋼中に含まれるものであるが、後述するＴｉやＮｂを添加した場合、ＴｉＮを形成して鋼の性質を高め、Ｎｂと結合して炭窒化物を形成して強度を増加させる。このため、Ｎ量として最低０．００１％必要である。しかしながら、Ｎ量の増加はＨＡＺ靭性、溶接性に極めて有害であり、本発明鋼においてはその上限は０．００６％である。
【００５９】
次に、必要に応じて含有することができるＮｉ、Ｃｕ、Ｃｒ、Ｔｉ、Ｖ、Ｃａ、ＲＥＭ、Ｍｇの添加理由と添加量範囲について説明する。
【００６０】
基本となる成分に、さらにこれらの元素を添加する主たる目的は、本発明鋼の優れた特徴を損なうことなく、強度、靭性などの特性を向上させるためである。したがって、その添加量は自ずと制限されるべき性質のものである。
【００６１】
Ｎｉは、溶接性、ＨＡＺ靭性に悪影響を及ぼすことなく母材の強度、靭性を向上させる。これら効果を発揮させるためには、少なくとも０．０５％以上の添加が必須である。一方、過剰な添加すると経済性を損なうだけでなく、溶接性に好ましくないため、上限を１．０％とした。
【００６２】
Ｃｕは、Ｎｉとほぼ同様の効果、現象を示し、上限の１．０％は溶接性劣化に加え、過剰な添加は熱間圧延時にＣｕ−クラックが発生し製造困難となるため規制される。下限は実質的な効果が得られるための最小量とすべきで０．０５％である。
【００６３】
Ｃｒは、母材の強度、靭性をともに向上させる。しかし、添加量が多すぎると母材、溶接部の靭性および溶接性を劣化させるため、限定範囲を０．０５〜１．０％とした。
【００６４】
上記、Ｃｕ、Ｎｉ、Ｃｒは、母材の強度、靭性上の観点のみならず、耐候性にも有効であり、そのような目的においては、溶接性を損ねない範囲で添加することが好ましい。
【００６５】
ＴｉもＮｂと同様に高温強度上昇に有効である。とくに、母材および溶接部靭性に対する要求が厳しい場合には、添加することが好ましい。なぜならばＴｉは、Ａｌ量が少ないとき（例えば０．００３％以下）、Ｏと結合してＴｉ2Ｏ3を主成分とする析出物を形成、粒内変態フェライト生成の核となり溶接部靭性を向上させる。また、ＴｉはＮと結合してＴｉＮとしてスラブ中に微細析出し、加熱時のγ粒の粗大化を抑え圧延組織の細粒化に有効であり、また鋼板中に存在する微細ＴｉＮは、溶接時に溶接熱影響部組織を細粒化するためである。これらの効果を得るためには、Ｔｉは最低０．００５％必要である。しかし多すぎるとＴｉＣを形成し、低温靭性や溶接性を劣化させるので、その上限は０．０２５％である。
【００６６】
Ｖは、Ｎｂとほぼ同様の作用を有するものであるが、Ｎｂに比べてその効果は小さい。また、Ｖは焼き入れ性にも影響を及ぼし、高温強度向上にも寄与する。Ｎｂと同様の効果は０．０１％未満では効果が少なく、上限は０．1％まで許容できる。
【００６７】
Ｃａ、ＲＥＭは不純物であるＳと結合し、靭性の向上や溶接部の拡散水素による誘起割れを抑制する働きを有するが、多すぎると粗大な介在物を形成し悪影響を及ぼすので、それぞれ０．０００５〜０．００４％、０．０００５〜０．００４％が適正範囲である。
【００６８】
Ｍｇは、溶接熱影響部においてオーステナイト粒の成長を抑制し、微細化する作用があり、溶接部の強靭化が図れる。このような効果を享受するためには、Ｍｇは０．０００１％以上必要である。一方、添加量が増えると添加量に対する効果代が小さくなり、経済性を失するため、上限は０．００６％とした。
【００６９】
鋼の個々の成分を限定しても、成分系全体が適切でないと優れた特性は得られない。このため、ＰCMの値を０．１８％以下の範囲に限定する。ＰCMは溶接性を表す指標で、低いほど溶接性は良好である。本発明鋼においては、ＰCMが０．１８％以下の範囲であれば優れた溶接性の確保が可能である。なお、溶接割れ感受性組成ＰCMは以下の式により定義する。
ＰCM＝Ｃ＋Ｓｉ／３０＋Ｍｎ／２０＋Ｃｕ／２０＋Ｎｉ／６０＋Ｃｒ／２０＋
Ｍｏ／１５＋Ｖ／１０＋５Ｂ
【００７０】
なお、Ｍｏ、Ｎｂ、Ｔｉ、Ｖと同様に、Ｗを適当量添加して、高温強度を確保することも本発明鋼の特性を向上させる有効な手段である。
【００７１】
さらに、鋼板の最終圧延方向の板厚断面方向１／４厚位置において、最終変態組織の旧オーステナイト粒径を平均円相当直径で１２０μｍ以下に限定する。これは、旧オーステナイト粒径が組織とともに靭性に大きな影響を及ぼすためで、特に本発明のような比較的多量のＭｏ添加鋼において靭性を高めるためには、旧オーステナイト粒径を小さく制御することは重要かつ必須である。前記旧オーステナイト粒径の限定理由は、発明者らの製造条件を種々変更した実験結果に基づくもので、平均円相当直径で１２０μｍ以下であれば、本発明よりも低Ｍｏである鋼と遜色ない靭性を確保できる。なお、旧オーステナイト粒は、その判別が必ずしも容易ではないケースも少なからずある。このような場合には、板厚１／４厚位置を中心として、鋼板の最終圧延方向と直角方向に採取した切り欠き付き衝撃試験片、例えば、ＪＩＳ Ｚ ２２０２ ４号試験片（２ｍｍＶノッチ）などを用い、十分低温で、脆性破壊させた際の破面単位を旧オーステナイト粒径と読み替え得る有効結晶粒径と定義し、その平均円相当直径を測定することとし、この場合でも同様に１２０μｍ以下であることが必要である。
【００７２】
【実施例】
転炉−連続鋳造−厚板工程で種々の鋼成分の鋼板（厚さ１５〜５０ｍｍ）を製造し、その強度、降伏比（ＹＲ）、靭性、７００℃、８００℃における降伏強さ、予熱なし（室温）におけるｙ割れ試験時のルート割れの有無等を調査した。
【００７３】
表１及び表２に比較鋼とともに本発明鋼の鋼成分を、表３に鋼板の製造条件および組織、表４に諸特性の調査結果を示す。
【００７４】
【表１】
【表２】
【表３】
【表４】
本発明鋼Ｎｏ．１〜１８の例では、全てミクロ組織がフェライト・ベイナイトの混合組織となっており、かつ旧オーステナイト粒径の平均円相当直径が１２０μｍ以下である。さらに、４９０ＭＰａ級鋼の常温の強度レベルを満足し、降伏比（ＹＲ）も７１〜７６％で８０％未満である。また、７００℃、８００℃のＹＳが常温での規格降伏強度のそれぞれ、６７％、２５％以上の良好な値で、実績高温常温降伏強度比（ｐ）についても、７００℃、８００℃でそれぞれ６４％、２３％以上の優れた値である。
【００７９】
これに対し、比較鋼Ｎｏ．１９では、Ｃが過剰であり、ベイナイト分率が過大となって、常温の降伏強度が４９０ＭＰａ級の上限を超える結果であった。また、高温強度については、絶対値としては高い値が得られているが、オーステナイトへの逆変態開始温度Ａｃ1 が８００℃以下となるため、７００℃における常温／高温の降伏強度比（ｐ）はｐ＜−０．００３３×Ｔ＋２．８０である。
【００８０】
比較鋼Ｎｏ．２０では、Ｃが不足であり、常温、高温ともに４９０ＭＰａ級として降伏強度が不足である。さらに、６００℃以上の高温における複合炭窒化相の生成量が５×１０-4未満であり、７００℃、８００℃における常温／高温の降伏強度比（ｐ）もｐ＜−０．００３３×Ｔ＋２．８０と低い。
【００８１】
比較鋼Ｎｏ．２１では、Ｍｎ量が０．５％を超えているため、常温での固溶強化効果が過剰となって、常温の降伏強度が４９０ＭＰａ級の規格値上限を超え、ＹＲも８０％超であった。また、Ａｃ１が８００℃未満となり、十分な整合／半整合析出強化効果が得られたかったため、７００℃、８００℃において、常温／高温降伏強度比（ｐ）がｐ＜−０．００３３×Ｔ＋２．８０である。
【００８２】
逆に、比較鋼Ｎｏ．２２では、Ｍｎ量が０．１％未満のため、常温での固溶強化効果が不足となって、常温及び７００℃の降伏強度、常温の引張り強度が４９０ＭＰａ級の規格値下限を下回った。
【００８３】
比較鋼Ｎｏ．２３では、Ｐが０．０２％を超えているため、母材の延性脆性遷移温度、０℃での再現ＨＡＺの吸収エネルギー値ともに劣化している。
【００８４】
比較鋼Ｎｏ．２４では、Ｓが０．０１％を超えているため、比較鋼Ｎｏ．２３と同様に、母材の延性脆性遷移温度、０℃での再現ＨＡＺの吸収エネルギー値ともに劣化している。
【００８５】
比較鋼Ｎｏ．２５ではＭｏの添加量不足により、炭窒化析出相、ＢＣＣ相中固溶Ｍｏがともに不足したため、常温強度、ＹＲ等は良好な結果であるが、７００℃の降伏強度が２１７ＭＰａ（４９０ＭＰａ級常温規格強度の２／３）未満で、８００℃の強度も６２ＭＰａ（４９０ＭＰａ級常温規格強度の２／９）未満と低い。また、実績の高温／常温降伏強度比（ｐ）も７００℃、８００℃について、それぞれ、４５％、１５％と低く、ｐ＜−０．００３３×Ｔ＋２．８０である。
【００８６】
比較鋼Ｎｏ．２６では、Ｍｏ量が過剰で、ミクロ組織がベイナイト単相となり、常温の降伏強度及び引張り強度が４９０ＭＰａ級の規格値上限を超えている。さらに、母材材質の不均一性が増大し、溶接割れ感受性組成ＰＣＭが０．１８％であるにも関わらず、予熱なしでのｙ割れ試験においてルート割れが発生した。また、再現ＨＡＺの吸収エネルギー値も低い。
【００８７】
比較鋼Ｎｏ．２７では、ＢＣＣ相中の固溶Ｍｏ、Ｎｂ量、複合炭窒化相の生成量ともに十分であり、高温／常温降伏強度比（ｐ）については、７００℃、８００℃ともにｐ≧−０．００３３×Ｔ＋２．８０を満足しているが、Ｎｂ量が不足したため、４９０ＭＰａ級として７００℃、８００℃の降伏強度が不足した。
【００８８】
逆に、比較鋼Ｎｏ．２８では、Ｎｂ量が過剰であるため、高温強度については高い値が得られるが、再現ＨＡＺの吸収エネルギー値は低い。
【００８９】
比較鋼Ｎｏ．２９では、γ粒が粗大であるため、再現ＨＡＺの吸収エネルギー値は低い。
【００９０】
比較鋼Ｎｏ．３０では、Ｔｉ量が過剰であるため、母材の延性脆性遷移温度、再現ＨＡＺ吸収エネルギー値ともに劣化している。
【００９１】
比較鋼Ｎｏ．３１では、Ｂ添加量が不足し、十分な焼入れ性を得ることができず、ミクロ組織のベイナイト分率が過少のため、常温、高温ともに降伏強度が４９０ＭＰａ級の規格値下限を下回った。
【００９２】
比較鋼Ｎｏ．３２では、Ｂ添加量が過剰なため、母材の延性脆性遷移温度は０℃近傍にあり、再現ＨＡＺの吸収エネルギー値は低い。
【００９３】
比較鋼Ｎｏ．３３では、Ａｌ量が０．０６％を超えているため、母材の延性脆性遷移温度は０℃近傍にあり、再現ＨＡＺ靭性も低い。
【００９４】
比較鋼Ｎｏ．３４では、Ｎ量が０．００６％を超えているため、再現ＨＡＺ靭性は低い。
【００９５】
比較鋼Ｎｏ．３５では、ＰCM値が０．１８％を超えており、予熱なしでのｙ割れ試験においてルート割れが発生した。また、再現ＨＡＺ吸収エネルギー値も低い。
【００９６】
比較鋼Ｎｏ．３６では、再加熱温度が１１００℃未満のため、再加熱時に添加合金元素がオーステナイト中に固溶せずに十分な析出強化が得られず、常温については降伏強度、引張り強度、ＹＲともに良好な結果であるが、７００℃の降伏強度が２１７ＭＰａ（４９０ＭＰａ級常温規格強度の２／３）未満で、８００℃の強度も７２ＭＰａ（４９０ＭＰａ級常温規格強度の２／９）未満と低い。さらに、実績の高温／常温降伏強度比（ｐ）は、７００℃、８００℃について、ｐ＜−０．００３３×Ｔ＋２．８０である。
【００９７】
比較鋼Ｎｏ．３７では、再加熱温度が１２５０℃を超えたため、再加熱時にオーステナイト粒が粗大化し、再現ＨＡＺの吸収エネルギー値が低くなっている。
【００９８】
比較鋼Ｎｏ．３８では、本願発明鋼Ｎｏ．１０と同成分であるが、１１００℃以下での累積圧下量が３０％未満のため、旧オーステナイト粒が粗大であり、再現ＨＡＺ靭性が低い。
【００９９】
比較鋼Ｎｏ．３９では、本願発明鋼Ｎｏ．１０と同成分であるが、８５０℃未満の温度で圧延を行ったため、Ｎｂ、Ｔｉ、Ｖの析出が促進され十分な析出強化が得られず、常温強度については４９０ＭＰａ級の規格値を満足するが、高温の降伏強度が不足し、実績の高温／常温降伏強度比（ｐ）は、７００℃、８００℃について、ｐ＜−０．００３３×Ｔ＋２．８０である。
【０１００】
比較鋼Ｎｏ．４０では、再加熱温度が１２５０℃と高いため、圧延終了後のオーステナイト粒が１２０μｍ超と粗大であり、フェライトの変態が抑制され、ベイナイト単相のミクロ組織となり、高温強度については高い値が得られているが、常温の降伏強度が４９０ＭＰａ級の規格上限を超過した。
【０１０１】
比較鋼Ｎｏ．４１では、圧延後水冷により冷却速度が過大となり、ベイナイト分率が過剰（＞９０％）となって、常温の降伏強度、引張り強度が４９０ＭＰａ級としての規格値上限を超え、ＹＲも８０％超であった。
【０１０２】
比較鋼Ｎｏ．４２では、圧延後水冷を行うことにより常温強度の上昇を図ったが、板厚が大きく１／４厚部におけるγ／α変態温度近傍での冷却速度が不足のため、フェライト分率が過大（＞８０％：ベイナイト分率＜２０％）となり、常温での固溶強化効果が不足となって、常温の引張り強度が規格値下限を下回り、７００℃、８００℃の降伏強度が、それぞれ、２１７ＭＰａ未満、７２ＭＰａ未満と低い。
【０１０３】
比較鋼Ｎｏ．４３では、板厚２５ｍｍ超であるため、加速冷却を適用し、０．３Ｋｓ-1以上の冷却速度の確保を図ったが、水冷開始温度が７００℃未満であり、圧延終了後〜冷却開始（６９０℃）の冷却速度が０．３Ｋｓ-1以下となり、水冷開始前にフェライトの変態が進行したため、ベイナイト分率が２０％未満となって、常温、高温ともに４９０ＭＰａ級として強度が不足した。
【０１０４】
比較鋼Ｎｏ．４４では、Ｔｉ量とＮ量がともに少なく、かつ、再加熱温度も１２５０℃と高いため、再加熱時にオーステナイトが１２０μｍ超に粗大化し、フェライトの変態が抑制され、ベイナイト単相のミクロ組織となり、高温強度については高い値が得られているが、常温の降伏強度、引張り強度が４９０ＭＰａ級の規格上限を超過した。
【０１０５】
【発明の効果】
本発明の化学成分及び製造法で製造した鋼材は、ミクロ組織がフェライト・ベイナイトの混合組織であり、常温強度が４９０ＭＰａの規格値を満足し、ＹＲが８０％以下、７００℃、８００℃の降伏強度がそれぞれ常温規格値の２／３以上、２／９以上等の特性を持ち、実績高温／常温降伏強度比（ｐ）が、７００℃〜８００℃において、ｐ≧−０．００３３×Ｔ＋２．８０を満足し、建築用耐火鋼材としての必要な特性を兼ね備えており、従来になく全く新しい鋼材である。[0001]
BACKGROUND OF THE INVENTION
The present invention is excellent in high-temperature strength in a relatively short time of about 1 hour in a temperature range of 600 ° C. or higher and 800 ° C. or lower used for general structures such as buildings, civil engineering, offshore structures, shipbuilding and storage tanks. The present invention relates to a method for producing high-strength steel (steel plate, steel pipe, section steel, wire rod) for building structures to which low alloy carbon is added.
[0002]
[Prior art]
For example, in the fields of construction and civil engineering, steel materials standardized by JIS and the like are widely used as various construction steel materials. In addition, since the strength of general steel for building structures is lowered from about 350 ° C., the allowable temperature is about 500 ° C.
[0003]
In other words, when using the steel materials described above for buildings such as buildings, offices, residences, and multistory parking lots, it is obliged to provide sufficient fireproof coating to ensure safety in fire. The laws and regulations stipulate that the temperature of the steel material does not exceed 350 ° C during a fire.
[0004]
This is because the steel material has a yield strength of about 2/3 of room temperature at about 350 ° C., which is lower than the required strength. When steel is used for a building, it is used with a fireproof coating so that the temperature of the steel does not reach 350 ° C. during a fire. For this reason, it is inevitable that the fireproof coating cost will be higher than the steel material cost, and the construction cost will increase significantly.
[0005]
In order to solve the above-mentioned problems, refractory steel having high temperature proof stress has been developed.
[0006]
In the case of steel having high-temperature strength at 600 ° C. or higher, it is generally called refractory steel, and refractory steel having a high-temperature strength of 2/3 or higher of normal temperature yield strength at 600 ° C. (see, for example, Patent Document 1), A refractory steel excellent in high-temperature strength at 700 ° C. (for example, see Patent Document 2) has been proposed. In other examples of the invention related to 600 ° C. refractory steel, the yield strength at 600 ° C. is generally 2/3 or more of the normal temperature yield strength.
[0007]
However, for 700 ° C. refractory steel and 800 ° C. refractory steel, there is no general rule for setting the high temperature strength (ratio to the room temperature yield strength) at present. For example, Patent Document 1 is a steel to which a considerable amount of Mo and Nb are added, and the proof stress at 600 ° C. ensures 70% or more of the normal temperature proof strength, but the proof strength at 700 ° C. to 800 ° C. is not shown. . Also, when the proof stress at 600 ° C is about 70% of the normal temperature proof strength, considering the temperature rise at the time of fire, the fireproof covering amount can be reduced, but the structures that can be omitted are multi-story parking lots, atriums, etc. Because it is limited to open spaces, its use in fire-proof coatings is significantly limited.
[0008]
Moreover, in patent document 2, although the proof stress of 700 degreeC ensures 56% or more of normal temperature proof stress by making a microstructure into bainite with the steel which added a considerable amount of Mo and Nb, the proof stress of 800 degreeC Is not shown.
[0009]
That is, steels having a high temperature strength of about 600 ° C. as in these examples have already been used in the market, and an invention of a steel material that ensures a certain strength at 700 ° C. has been made. Stable production of practical steel that can ensure high temperature strength at ℃ has been difficult.
[0010]
[Patent Document 1]
Japanese Patent Laid-Open No. 2-77523
[Patent Document 2]
Japanese Patent Laid-Open No. 10-68044
[0011]
[Problems to be solved by the invention]
As described above, when steel is used for a building, normal steel has low high-temperature strength, so it cannot be used for uncoated or fire-resistant coating reduction, and expensive fire-resistant coating has to be applied.
[0012]
In addition, even with newly developed steel, the fire resistance temperature is limited to a guarantee of 600 to 700 ° C, and the development of a steel material that enables the use of a fireproof coating at 700 ° C to 800 ° C and the omission of the fireproof coating process thereby. Was desired.
[0013]
An object of the present invention is to provide a high-strength steel excellent in high-temperature strength and weldability in a temperature range of 600 ° C. or higher and 800 ° C. or lower, and a production method capable of supplying the steel stably industrially.
[0014]
[Means for Solving the Problems]
In order to overcome the above-mentioned problems, the present invention achieves the object by setting the microstructure and the amount of additive alloy elements in the optimum range, and the gist thereof is as follows.
[0015]
(1) The steel component is mass%,
C: 0.005% or more and less than 0.04%,
Si: 0.5% or less,
Mn: 0.1% or more and 0.5% or less,
P: 0.02% or less,
S: 0.01% or less,
Mo: 0.3 to 1.5%,
Nb: 0.03-0.15%,
B: 0.0005 to 0.003%,
Al: 0.06% or less,
N: 0.006% or less,
A 490 MPa class high-strength steel excellent in high-temperature strength, characterized in that the balance consists of iron and inevitable impurities.
[0016]
(2) In mass%,
Ni: 0.05 to 1.0%,
Cu: 0.05 to 1.0%,
Cr: 0.05 to 1.0%,
Ti: 0.005 to 0.025%,
V: 0.01 to 0.1%
The 490 MPa class high-tensile steel excellent in high-temperature strength as described in (1) above, which contains 1 type or 2 types or more in the range described above.
[0017]
(3) In mass%,
Ca: 0.0005 to 0.004%,
REM: 0.0005 to 0.004%,
Mg: 0.0001 to 0.006%
The 490 MPa class high-tensile steel excellent in high-temperature strength as described in (1) or (2) above, which contains any one or more of the above.
[0018]
(4) The high temperature normal temperature yield stress ratio p (= high temperature yield stress / normal temperature yield stress) obtained by making the yield stress at high temperature dimensionless by the normal temperature yield stress is a steel material temperature T (° C.) of 600 ° C. or higher and 800 ° C. or lower. 490 MPa class high strength steel excellent in high temperature strength according to any one of the above (1) to (3), wherein p ≧ −0.0033 × T + 2.80 is satisfied.
[0019]
(5) It is a mixed structure of ferrite and bainite at normal temperature, and the temperature (Ac1) that reversely transforms to austenite when heated at a high temperature equivalent to a fire is over 800 ° C. 490 MPa class high strength steel excellent in high temperature strength described in any of the above.
[0020]
(6) A carbonitrided precipitation phase that is thermodynamically stable at high temperature in a mixed matrix structure of ferrite and bainite at a molar fraction of 5 × 10- 4While maintaining the above, the total amount of Mo and Nb dissolved in the BCC phase is 2 × 10 in terms of molar concentration.- 3The 490 MPa class high-tensile steel excellent in high-temperature strength as described in (5) above.
[0021]
(7) A carbonitride precipitation phase that is thermodynamically stable at a high temperature in a mixed matrix structure of ferrite and bainite at a molar fraction of 5 × 10-FourWhile maintaining the above, the total amount of Mo, Nb, V and Ti dissolved in the BCC phase is 2 × 10 in terms of molar concentration.-3The 490 MPa class high-tensile steel excellent in high-temperature strength as described in (5) above.
[0022]
(8) A 490 MPa class high tension excellent in high-temperature strength according to any one of (1) to (7) above, wherein the mixed structure of ferrite and bainite is a bainite fraction of 20 to 90%. steel.
[0023]
(9) The steel component according to any one of (1) to (3) above,
Pcm= C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 +
Mo / 15 + V / 10 + 5B
Weld cracking susceptibility composition P defined ascm490 MPa class high-tensile steel excellent in high-temperature strength and weldability, characterized in that is 0.18% or less.
[0024]
(10) In addition to the characteristics described in any one of (4) to (7) above, the bainite fraction is 20 to 90% as a mixed structure of ferrite and bainite. 490 MPa class high strength steel excellent in high temperature strength and weldability as described in (9) above.
[0025]
(11) The 490 MPa class high-tensile steel excellent in high-temperature strength and weldability as described in (9) or (10) above, wherein the average equivalent circle diameter of the prior austenite grains is 120 μm or less.
[0026]
(12) After reheating the steel slab or slab comprising the steel component according to any one of (1) to (3) above to a temperature range of 1100 to 1250 ° C, the cumulative reduction amount at 1100 ° C or less is set. A method for producing a high-strength steel of 490 MPa class excellent in high-temperature strength, characterized in that it is rolled at a temperature of 850 ° C. or higher at 30% or higher and the microstructure is a mixed structure of ferrite and bainite.
[0027]
(13) After reheating the steel slab or slab comprising the steel component according to any one of (1) to (3) above to a temperature range of 1100 to 1250 ° C, the cumulative reduction amount at 1100 ° C or less is set. 30% or more, rolling at a temperature of 850 ° C. or higher, and after cooling, the cooling rate from 800 ° C. or higher to 650 ° C. or lower is 0.3 Ks.-1As described above, the method for producing a 490 MPa class high strength steel excellent in high temperature strength according to any one of the above (4) to (7), wherein the microstructure is a mixed structure of ferrite and bainite.
[0028]
(14) AboveAny one of (1) to (3)After reheating the slab comprising the steel components described in 1 to a temperature range of 1100 to 1250 ° C., the cumulative reduction at 1100 ° C. or lower is set to 30% or more, and the steel is rolled at a temperature of 850 ° C. or higher. A method for producing a 490 MPa class high-tensile steel excellent in high-temperature strength and weldability, characterized by having a mixed structure of bainite.
[0029]
(15) AboveAny one of (1) to (3)After reheating the slab made of the steel component described in 1100 to 1250 ° C in a temperature range, the cumulative reduction at 1100 ° C or less is set to 30% or more and rolled at a temperature of 850 ° C or more. The cooling rate from the above temperature to 650 ° C. or less is 0.3 Ks.-1As described above, the microstructure is a mixed structure of ferrite and bainite.HighA method for producing a low-yield ratio 490 MPa class high-strength steel excellent in temperature strength and weldability.
[0030]
DETAILED DESCRIPTION OF THE INVENTION
Details of the present invention will be described below.
[0031]
The present inventors have already found steel having excellent high-temperature strength at 600 ° C. and 700 ° C. Steel with excellent high-temperature strength at 600 ° C. is already used in the construction field, but there is a very strong demand for steel that can withstand higher temperatures in the market.
[0032]
To increase the strength at high temperatures, the combined addition of Mo and Nb promotes the precipitation of carbonitrides that are stable at high temperatures, the dislocation density increases due to the bainite of the microstructure, and further due to the solid solution of Mo and Nb. It is effective to delay the dislocation recovery. However, if the fraction of hard bainite is excessive, the strength at normal temperature becomes excessive, so that the microstructure is a mixed structure of ferrite and bainite having an appropriate bainite fraction according to the required normal temperature strength. Low C is effective in building an appropriate microstructure and achieving the required normal temperature strength range. Lowering C also has the effect of increasing the thermodynamic stability of the mixed matrix structure of ferrite and bainite at high temperatures and increasing the reverse transformation temperature (Ac1) to austenite. However, in this case, it has been found that the microstructure and material are easily affected by the rolling conditions and the subsequent cooling conditions, and stable production is difficult. As a result of tackling the microstructure control and increasing the high-temperature strength, the inventors have found that an appropriate amount of B addition is effective for stabilization of production, and have reached the present invention.
[0033]
As a general welded structural steel, characteristics such as weldability and low yield strength ratio need to be provided in the same way as in the past, and therefore steel with excellent high-temperature strength at 700 ° C to 800 ° C is an extremely difficult task. there were. In order to solve this problem, the present inventors diligently studied, and the high temperature strengths of 700 ° C. and 800 ° C. are caused by precipitation strengthening by complex addition of alloy elements such as Mo, Nb, V, and Ti, and dislocation by microstructure bainite. It was found that the increase in density and further the delay of dislocation recovery due to solute Mo, Nb, V, and Ti are effective, and that Ti also has some effect. Therefore, in order to ensure all of the strength of 700 ° C. to 800 ° C., the strength at normal temperature, and the strength ratio between normal temperature and high temperature (YS ratio = high temperature strength / normal temperature strength) at the same time, the microstructure should be mixed with appropriate ferrite and bainite. It was found that it is important to obtain the microstructure, the thermal stability of the matrix structure at high temperatures, the appropriate coherent precipitation strengthening effect, and the dislocation recovery delay effect with the added alloy element amount in the optimum range.
[0034]
The yield strength of steel generally decreases rapidly from around 450 ° C. This is because the thermal activation energy decreases as the temperature rises, and the resistance that was effective at low temperatures against dislocation slip motion becomes ineffective. Because. The present inventors have found that a composite carbonitride of Mo, Nb, V, and Ti acts as an effective resistance up to a high temperature of about 600 ° C. with respect to the slip motion of dislocations. Furthermore, Mo, Nb, V, and Ti dissolved in the BCC phase are effective for dislocation recovery delay, and have the effect of increasing the temperature at which the sudden decrease in yield strength begins. It was. Therefore, at 700 ° C. to 800 ° C., the steel material temperature is T (° C.), and the high temperature normal temperature yield stress ratio p (= high temperature yield stress / normal temperature yield stress) satisfies p ≧ −0.0033 × T + 2.80. That is, in order for the yield stress ratio to be 49% and 16% or more, respectively, the composite carbonitrides of Mo, Nb, V, and Ti contained as constituent elements in the steel at the temperature are 5 in mole fraction. × 10-FourIn addition to the above, the total amount of Mo, Nb, V, and Ti, which are component elements dissolved in the BCC phase, is 2 × 10 in terms of molar concentration.-3It must be more than that.
[0035]
The composition of the composite carbonitride precipitation phase, which is important for high-temperature strength development, can be easily identified, for example, by analysis with an electron microscope or EDX.
[0036]
Also, the equilibrium formation amount of the thermodynamically stable precipitated phase and the solid solution alloy element amount in the BCC phase can be easily calculated from the additive alloy element amount by using commercially available thermodynamic calculation database software. is there.
[0037]
When the fraction of ferrite in the microstructure increases and the fraction of bainite decreases to less than 20%, the strength at normal temperature and high temperature decreases, and it is necessary to add more alloying elements such as Mo, Nb, Ti, and V. Arise. If the fraction of ferrite occupying the microstructure is excessive, it is difficult to ensure the strength at normal temperature and high temperature due to the increase of the additive alloy elements. Conversely, when the fraction of ferrite in the microstructure decreases and the fraction of bainite increases, the strength at room temperature and high temperature increases, and the amount of alloying elements such as Mo, Nb, Ti, and V needs to be reduced. Furthermore, if the bainite fraction in the microstructure exceeds 90% and is excessive, high temperature strength can be achieved, but the increase in room temperature strength, HAZ toughness, and weldability are significant and the amount of added alloy elements is reduced. It becomes difficult to ensure the required strength range, HAZ toughness and weldability.
[0038]
For this reason, in the steel of the present invention, the microstructure is a mixed structure of ferrite and bainite, and the fraction of bainite is in the range of 20% to 90%, but is preferably in the range of 30 to 70%.
[0039]
The present inventors have studied a method for maintaining the microstructure to be a mixed structure of ferrite and bainite and keeping the bainite fraction stably in the range of 20 to 90%, and found that addition of an appropriate amount of B is essential. It was.
[0040]
The reason why the present invention limits the steel components and the production method as described in the claims will be described.
[0041]
In order to ensure the strength at normal temperature and high temperature at the same time, it is necessary to add a considerable amount of alloying elements. In 490 MPa class high-tensile steel, Mo: 0.3 to 1.5%, Nb: 0.03 to 0 .15% is required.
[0042]
Further, addition of Ti: 0.005 to 0.025% and V: 0.01 to 0.1% is effective for improving the high temperature strength.
[0043]
Mo, Nb, Ti, V, etc. are mainly for securing high-temperature strength, and limiting the range of Si and Mn is for suppressing the normal-temperature strength within a predetermined range.
[0044]
The heating temperature of the steel is preferably a high temperature in order to make Mo, Nb, Ti, and V as solid as possible, but is limited to 1100 to 1250 ° C. from the viewpoint of securing the toughness of the base material.
[0045]
The rolling end temperature is 1100 ° C. or lower and 850 ° C. or higher because Mo, Nb, Ti, and V precipitate as carbides under a low temperature pressure, and 850 ° C. is the lower limit temperature, and the temperature exceeds 1100 ° C. This is because the toughness is insufficient when rolling is completed.
[0046]
In addition, even if it reheats to the temperature below Ac1 transformation point for the purpose, such as dehydrogenation, after manufacturing this invention steel, the characteristic of this invention steel is not impaired at all.
[0047]
Next, other component elements related to the present description and the addition amount thereof will be described.
[0048]
C has the most remarkable effect on the properties of the steel material and must be controlled in a narrow range. 0.005 or more and less than 0.04% is a limited range. If the amount of C is less than 0.005%, the strength is insufficient, and if it is 0.04% or more, if the cooling rate after rolling is excessive, the fraction of bainite increases, the strength exceeds, and conversely the cooling rate is too low. In the case of, the yield of bainite is lowered and the strength is insufficient. Furthermore, during high-temperature heating equivalent to fire, the mixed matrix structure of ferrite and bainite is kept thermodynamically stable, maintaining consistency with the composite carbonitride precipitates of Mo, Nb, V, and Ti. In order to secure C, C must be less than 0.04%.
[0049]
Si is an element contained in the deoxidized upper steel, and is effective for improving the strength of the base material at room temperature because it has a substitutional solid solution strengthening action. In particular, the effect of improving the high temperature strength above 600 ° C. is Absent. Moreover, since weldability and HAZ toughness will deteriorate when adding much, the upper limit was limited to 0.5%. Deoxidation of steel can be performed only with Ti and Al, and is preferably as low as possible from the viewpoints of HAZ toughness, hardenability, and the like, and it is not always necessary to add.
[0050]
Mn is an indispensable element for ensuring strength and toughness. However, Mn, which is a substitutional solid solution strengthening element, is effective for increasing the strength at room temperature. There is not much improvement effect. Therefore, in a steel containing a relatively large amount of Mo as in the present invention, the weldability is improved.cmFrom the viewpoint of reduction, it is necessary to be 0.9% or less. Furthermore, for the 490 MPa class high-strength steel for construction, the upper limit of the normal temperature strength was taken into consideration and the content was limited to 0.5% or less. By keeping the upper limit of Mn low, it is advantageous from the viewpoint of center segregation of the continuously cast slab. In addition, about a minimum, 0.1% or more of addition is required on the intensity | strength and toughness adjustment of a base material. Therefore, Mn is set to a range of 0.1 to 0.5%.
[0051]
In order to make the yield strength and tensile strength at room temperature within the required range of 490 MPa class high-tensile steel, the cooling rate from 800 ° C. to 650 ° C. after rolling is 0.3 Ks.-1It is necessary to do it above. That is, a relatively thin steel plate of less than about 25 mm can be manufactured by an air cooling or accelerated cooling (water cooling) process, and a relatively thick steel plate of about 25 mm or more can be manufactured by applying an accelerated cooling (water cooling) process.
[0052]
P is an impurity in the steel of the present invention, and a reduction in the amount of P tends to reduce the grain boundary fracture in the HAZ, so the smaller the better. If the content is large, the low temperature toughness of the base metal and the welded portion is deteriorated, so the upper limit was made 0.02%.
[0053]
S, like P, is an impurity in the steel of the present invention, and is preferably as small as possible from the viewpoint of the low temperature toughness of the base material. If the content is large, the low temperature toughness of the base metal and the welded portion is deteriorated, so the upper limit was made 0.01%.
[0054]
Mo is an indispensable element for securing high-temperature strength at 700 ° C. and 800 ° C., and is one of the most important elements in the present invention. If considering only high-temperature strength, the lower limit can be relaxed, but from the viewpoint of lowering the yield ratio described later, two-phase heat treatment of ferrite and austenite, and then tempering as necessary, In order to ensure high strength and high toughness at room temperature, the lower limit was made 0.3%. On the other hand, addition of more than 1.5% makes it difficult to control the base material (variation control and toughness deterioration) and loses economic efficiency, so 0.3 to 1.5% is limited. is there.
[0055]
Nb is an element that plays an important role in securing high temperature strength at 700 ° C. and 800 ° C. in the present invention in which a relatively large amount of Mo is added. First, as a general effect, it is an element useful for raising the recrystallization temperature of austenite and maximizing the effect of controlled rolling during hot rolling. It also contributes to re-heating prior to rolling, normalizing, and refinement of heated austenite during quenching. Furthermore, it has the effect of improving the high temperature strength by precipitation strengthening and suppressing dislocation recovery, and contributes to the improvement of high temperature strength by the combined addition with Mo. If it is less than 0.03%, the effect of suppressing precipitation hardening and dislocation recovery at 700 ° C. and 800 ° C. is small, and if it exceeds 0.15%, the degree of effect decreases with respect to the added amount, which is not economically preferable. Moreover, the toughness at the time of welding also falls. Therefore, 0.03 to 0.15% is the limited range.
[0056]
B is extremely important in controlling the strength through the production fraction of bainite. In other words, B segregates at the austenite grain boundaries and suppresses the formation of ferrite, thereby improving the hardenability and stably generating bainite even when the cooling rate is relatively low such as air cooling. It is valid. In order to enjoy this effect, at least 0.0005% or more is necessary. However, too much addition not only saturates the effect of improving hardenability, but also may form B precipitates that are detrimental to embrittlement and toughness of prior austenite grain boundaries, so the upper limit is 0.003%. did. In cases where stress corrosion cracking is a concern, such as for tank steel, reduction of the hardness of the base metal and the weld heat affected zone is often the point (for example, prevention of sulfide stress corrosion cracking (SCC)). Therefore, HRC ≦ 22 (HV ≦ 248) is essential), and in such a case, excessive B addition that increases hardenability is not preferable.
[0057]
Al is an element generally contained in deoxidized steel, but Si or Ti is sufficient for deoxidation, and the lower limit is not limited (including 0%) in the steel of the present invention. However, when the amount of Al increases, not only the cleanliness of the steel deteriorates but also the toughness of the weld metal deteriorates, so the upper limit was made 0.06%.
[0058]
N is contained in the steel as an unavoidable impurity. However, when Ti or Nb, which will be described later, is added, TiN is formed to enhance the properties of the steel, and it combines with Nb to form a carbonitride. Increase strength. For this reason, the N amount is required to be at least 0.001%. However, the increase in the amount of N is extremely harmful to the HAZ toughness and weldability, and the upper limit of the steel of the present invention is 0.006%.
[0059]
Next, the reason for addition of Ni, Cu, Cr, Ti, V, Ca, REM, and Mg and the range of the amount that can be contained as necessary will be described.
[0060]
The main purpose of adding these elements to the basic components is to improve properties such as strength and toughness without impairing the excellent characteristics of the steel of the present invention. Therefore, the amount of addition is naturally limited.
[0061]
Ni improves the strength and toughness of the base material without adversely affecting weldability and HAZ toughness. In order to exert these effects, addition of at least 0.05% is essential. On the other hand, excessive addition not only impairs economic efficiency but also is unfavorable for weldability, so the upper limit was made 1.0%.
[0062]
Cu exhibits substantially the same effects and phenomena as Ni, and the upper limit of 1.0% is restricted because weldability deteriorates, and excessive addition causes Cu-cracks during hot rolling, which makes manufacturing difficult. The lower limit should be the minimum amount for obtaining a substantial effect, and is 0.05%.
[0063]
Cr improves both the strength and toughness of the base material. However, if the addition amount is too large, the base material, the toughness of the welded portion and the weldability are deteriorated, so the limited range was made 0.05 to 1.0%.
[0064]
Cu, Ni, and Cr are effective not only in terms of the strength and toughness of the base material but also in weather resistance. For such purposes, it is preferable to add Cu, Ni, and Cr in a range that does not impair the weldability.
[0065]
Ti, as well as Nb, is effective for increasing the high temperature strength. In particular, when the requirements for the base material and weld toughness are severe, it is preferable to add them. This is because Ti combines with O when Ti content is small (for example, 0.003% or less).2OThreeIs formed as a main component, and becomes the nucleus of the formation of intragranular transformation ferrite, and improves the toughness of the weld. Ti is combined with N and finely precipitated in the slab as TiN, which suppresses the coarsening of γ grains during heating and is effective for refining the rolled structure. The fine TiN present in the steel sheet is welded. This is to sometimes refine the weld heat affected zone structure. In order to obtain these effects, Ti needs to be at least 0.005%. However, if it is too much, TiC is formed and the low temperature toughness and weldability are deteriorated, so the upper limit is 0.025%.
[0066]
V has substantially the same action as Nb, but its effect is smaller than that of Nb. V also affects the hardenability and contributes to the improvement of high temperature strength. The effect similar to Nb is less if it is less than 0.01%, and the upper limit is acceptable up to 0.1%.
[0067]
Ca and REM combine with S, which is an impurity, to improve toughness and suppress induced cracking caused by diffusion hydrogen in the weld. However, if too much, coarse inclusions are formed and adversely affected. 0005 to 0.004% and 0.0005 to 0.004% are appropriate ranges.
[0068]
Mg suppresses the growth of austenite grains in the weld heat-affected zone and has the effect of miniaturization, so that the weld zone can be strengthened. In order to enjoy such effects, Mg needs to be 0.0001% or more. On the other hand, as the amount added increases, the effect on the amount added decreases and the economy is lost, so the upper limit was made 0.006%.
[0069]
Even if the individual components of the steel are limited, excellent properties cannot be obtained unless the entire component system is appropriate. For this reason, PcmIs limited to a range of 0.18% or less. PcmIs an index representing weldability, and the lower, the better the weldability. In the steel of the present invention, PcmIf it is 0.18% or less of range, it is possible to ensure excellent weldability. In addition, weld crack sensitivity composition PcmIs defined by the following equation.
Pcm= C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 +
Mo / 15 + V / 10 + 5B
[0070]
As in the case of Mo, Nb, Ti, and V, adding an appropriate amount of W to ensure high temperature strength is also an effective means for improving the properties of the steel of the present invention.
[0071]
Furthermore, the old austenite grain size of the final transformation structure is limited to 120 μm or less in terms of the average equivalent circle diameter at the position of ¼ thickness in the sheet thickness cross-section direction in the final rolling direction of the steel sheet. This is because the prior austenite grain size has a great influence on the toughness as well as the structure. In particular, in order to increase the toughness in a relatively large amount of Mo-added steel as in the present invention, it is necessary to control the prior austenite grain size to be small. Important and essential. The reason for limiting the prior austenite grain size is based on the experimental results obtained by changing the production conditions of the inventors. If the average equivalent circle diameter is 120 μm or less, it is comparable to steel having a lower Mo than the present invention. Toughness can be secured. In addition, there are not a few cases where it is not always easy to distinguish old austenite grains. In such a case, an impact test piece with a notch sampled in a direction perpendicular to the final rolling direction of the steel sheet centering on the position where the plate thickness is 1/4 thickness, for example, a JIS Z 2204 No. 4 test piece (2 mmV notch), etc. Is defined as the effective grain size that can be read as the prior austenite grain size at a sufficiently low temperature, and the average equivalent circle diameter is measured. It is necessary to be.
[0072]
【Example】
Manufacture steel plates (thickness 15-50mm) of various steel components in the converter-continuous casting-thick plate process, strength, yield ratio (YR), toughness, yield strength at 700 ° C and 800 ° C, no preheating The presence or absence of root cracks during the y-crack test at room temperature was investigated.
[0073]
Tables 1 and 2 show the steel components of the steel of the present invention together with the comparative steel, Table 3 shows the manufacturing conditions and structure of the steel sheet, and Table 4 shows the investigation results of various properties.
[0074]
[Table 1]
[Table 2]
[Table 3]
[Table 4]
Invention Steel No. In the examples 1 to 18, all the microstructures are a mixed structure of ferrite and bainite, and the average equivalent-circle diameter of the prior austenite grain size is 120 μm or less. Furthermore, the strength level of normal temperature of 490 MPa class steel is satisfied, and the yield ratio (YR) is 71-76% and less than 80%. In addition, YS at 700 ° C and 800 ° C is a good value of 67% and 25% or more of the standard yield strength at room temperature, respectively, and the actual high temperature room temperature yield strength ratio (p) is also at 700 ° C and 800 ° C, respectively. Excellent values of 64% and 23% or more.
[0079]
On the other hand, comparative steel No. In No. 19, C was excessive, the bainite fraction was excessive, and the yield strength at room temperature exceeded the upper limit of the 490 MPa class. Further, as for the high temperature strength, a high value is obtained as an absolute value, but the reverse transformation start temperature Ac to austenite Ac1 Is 800 ° C. or less, the normal temperature / high temperature yield strength ratio (p) at 700 ° C. is p <−0.0033 × T + 2.80.
[0080]
Comparative steel No. In No. 20, C is insufficient, and yield strength is insufficient as a 490 MPa class at both normal temperature and high temperature. Furthermore, the amount of composite carbonitriding phase produced at a high temperature of 600 ° C. or higher is 5 × 10-FourThe yield strength ratio (p) at room temperature / high temperature at 700 ° C. and 800 ° C. is also low, p <−0.0033 × T + 2.80.
[0081]
Comparative steel No. In No. 21, since the Mn content exceeds 0.5%, the solid solution strengthening effect at room temperature becomes excessive, the yield strength at room temperature exceeds the upper limit of the standard value of the 490 MPa class, and the YR exceeds 80%. It was. Ac1Was less than 800 ° C., and a sufficient matched / semi-matched precipitation strengthening effect was not obtained, so that the normal temperature / high temperature yield strength ratio (p) is p <−0.0033 × T + 2.80 at 700 ° C. and 800 ° C. .
[0082]
On the contrary, comparative steel No. In No. 22, since the amount of Mn was less than 0.1%, the effect of solid solution strengthening at room temperature was insufficient, and the yield strength at room temperature and 700 ° C. and the tensile strength at room temperature were below the lower limit of the standard value of the 490 MPa class.
[0083]
Comparative steel No. 23, since P exceeds 0.02%, both the ductile brittle transition temperature of the base material and the absorption energy value of the reproduced HAZ at 0 ° C. are deteriorated.
[0084]
Comparative steel No. 24, since S exceeds 0.01%, comparative steel No. 23, both the ductile brittle transition temperature of the base material and the reconstructed HAZ absorbed energy value at 0 ° C. are deteriorated.
[0085]
Comparative steel No. In No. 25, both the carbonitride precipitation phase and the solid solution Mo in the BCC phase were insufficient due to insufficient addition of Mo, so that the room temperature strength, YR, etc. are good results, but the yield strength at 700 ° C. is 217 MPa (490 MPa class room temperature standard). The strength at 800 ° C. is less than 62 MPa (2/9 of the 490 MPa class normal temperature standard strength). Moreover, the actual high temperature / normal temperature yield strength ratio (p) is as low as 45% and 15% at 700 ° C. and 800 ° C., respectively, p <−0.0033 × T + 2.80.
[0086]
Comparative steel No. In No. 26, the amount of Mo is excessive, the microstructure becomes a bainite single phase, and the yield strength and tensile strength at room temperature exceed the upper limit of the standard value of the 490 MPa class. Furthermore, the non-uniformity of the base material is increased and the weld cracking susceptibility composition PCMDespite being 0.18%, root cracks occurred in the y crack test without preheating. Also, the absorption energy value of the reproduced HAZ is low.
[0087]
Comparative steel No. 27, the amount of solute Mo and Nb in the BCC phase and the amount of the composite carbonitriding phase are sufficient, and the high temperature / normal temperature yield strength ratio (p) is p ≧ −0.0033 for both 700 ° C. and 800 ° C. XT + 2.80 was satisfied, but because the Nb content was insufficient, the yield strength at 700 ° C. and 800 ° C. as the 490 MPa class was insufficient.
[0088]
On the contrary, comparative steel No. In No. 28, since the amount of Nb is excessive, a high value is obtained for the high-temperature strength, but the absorption energy value of the reproduced HAZ is low.
[0089]
Comparative steel No. In 29, since the γ grains are coarse, the absorption energy value of the reproduced HAZ is low.
[0090]
Comparative steel No. In No. 30, since the Ti amount is excessive, both the ductile brittle transition temperature and the reproduced HAZ absorbed energy value of the base material are deteriorated.
[0091]
Comparative steel No. In No. 31, the amount of B added was insufficient, sufficient hardenability could not be obtained, and the bainite fraction of the microstructure was too small, so that the yield strength was below the lower limit of the standard value of the 490 MPa class at both room temperature and high temperature.
[0092]
Comparative steel No. In No. 32, since the addition amount of B is excessive, the ductile brittle transition temperature of the base material is in the vicinity of 0 ° C., and the absorption energy value of the reproduced HAZ is low.
[0093]
Comparative steel No. In No. 33, since the Al content exceeds 0.06%, the ductile brittle transition temperature of the base material is in the vicinity of 0 ° C., and the reproduced HAZ toughness is also low.
[0094]
Comparative steel No. In No. 34, since the N content exceeds 0.006%, the reproduced HAZ toughness is low.
[0095]
Comparative steel No. In 35, PcmThe value exceeded 0.18%, and root cracking occurred in the y cracking test without preheating. Also, the reproduced HAZ absorbed energy value is low.
[0096]
Comparative steel No. In No. 36, since the reheating temperature is less than 1100 ° C., the additive alloy element does not dissolve in austenite at the time of reheating, and sufficient precipitation strengthening cannot be obtained, and the yield strength, tensile strength, and YR are good at room temperature. As a result, the yield strength at 700 ° C. is less than 217 MPa (2/3 of the 490 MPa class room temperature standard strength) and the strength at 800 ° C. is also less than 72 MPa (2/9 of the 490 MPa class room temperature standard strength). Furthermore, the actual high temperature / normal temperature yield strength ratio (p) is p <−0.0033 × T + 2.80 for 700 ° C. and 800 ° C.
[0097]
Comparative steel No. In No. 37, since the reheating temperature exceeded 1250 ° C., the austenite grains became coarse during the reheating, and the absorption energy value of the reproduced HAZ was low.
[0098]
Comparative steel No. 38, the present invention steel No. 10 is the same component, but since the cumulative reduction at 1100 ° C. or less is less than 30%, the prior austenite grains are coarse and the reproduced HAZ toughness is low.
[0099]
Comparative steel No. 39, the present invention steel No. 10 is the same component, but since rolling was performed at a temperature lower than 850 ° C., precipitation of Nb, Ti, V was promoted and sufficient precipitation strengthening was not obtained, and the normal temperature strength satisfied the standard value of 490 MPa class. However, the yield strength at high temperatures / room temperature yield strength ratio (p) is p <−0.0033 × T + 2.80 at 700 ° C. and 800 ° C.
[0100]
Comparative steel No. In No. 40, since the reheating temperature is as high as 1250 ° C., the austenite grains after rolling are as coarse as over 120 μm, the transformation of ferrite is suppressed, a microstructure of bainite single phase is obtained, and a high value is obtained for high temperature strength. However, the yield strength at room temperature exceeded the upper limit of the standard of 490 MPa class.
[0101]
Comparative steel No. In No. 41, the cooling rate becomes excessive due to water cooling after rolling, the bainite fraction becomes excessive (> 90%), the yield strength at normal temperature and the tensile strength exceed the upper limit of the standard value as 490 MPa class, and YR exceeds 80%. Met.
[0102]
Comparative steel No. In No. 42, the strength at room temperature was increased by performing water cooling after rolling, but the ferrite fraction was excessive because the plate thickness was large and the cooling rate in the vicinity of the γ / α transformation temperature in the ¼ thickness portion was insufficient. > 80%: Bainitic fraction <20%), the effect of solid solution strengthening at room temperature is insufficient, the tensile strength at room temperature is below the lower limit of the standard value, and the yield strength at 700 ° C. and 800 ° C. is 217 MPa, respectively. Less than or less than 72 MPa.
[0103]
Comparative steel No. In 43, since the plate thickness exceeds 25 mm, accelerated cooling is applied and 0.3 Ks is applied.-1Although the above cooling rate was ensured, the water cooling start temperature was less than 700 ° C., and the cooling rate after rolling to the start of cooling (690 ° C.) was 0.3 Ks.-1Since the ferrite transformation proceeded before the start of water cooling, the bainite fraction was less than 20%, and the room temperature and high temperature were both 490 MPa class and the strength was insufficient.
[0104]
Comparative steel No. In No. 44, both the Ti amount and the N amount are small, and the reheating temperature is as high as 1250 ° C., so the austenite is coarsened to over 120 μm during reheating, the ferrite transformation is suppressed, and a microstructure of a bainite single phase is formed. Although high values were obtained for the high temperature strength, the yield strength and tensile strength at room temperature exceeded the upper limit of the standard of 490 MPa class.
[0105]
【The invention's effect】
The steel material produced by the chemical composition and production method of the present invention has a microstructure of a mixed structure of ferrite and bainite, satisfies the standard value of normal temperature strength of 490 MPa, YR is 80% or less, yield of 700 ° C. and 800 ° C. The strength is 2/3 or more of the normal temperature standard value, 2/9 or more, etc., and the actual high temperature / normal temperature yield strength ratio (p) is 700 ° C. to 800 ° C., p ≧ −0.0033 × T + 2. It is a completely new steel material that satisfies the requirements of 80 and has the necessary characteristics as a fireproof steel material for construction.
Claims (15)
Ｃ：０．００５％以上０．０４％未満、
Ｓｉ：０．５％以下、
Ｍｎ：０．１％以上０．５％以下、
Ｐ：０．０２％以下、
Ｓ：０．０１％以下、
Ｍｏ：０．３〜１．５％、
Ｎｂ：０．０３〜０．１５％、
Ｂ：０．０００５〜０．００３％、
Ａｌ：０．０６％以下、
Ｎ：０．００６％以下、
かつ、残部が鉄および不可避的不純物からなることを特徴とする高温強度に優れた４９０ＭＰａ級高張力鋼。Steel component is mass%,
C: 0.005% or more and less than 0.04%,
Si: 0.5% or less,
Mn: 0.1% or more and 0.5% or less,
P: 0.02% or less,
S: 0.01% or less,
Mo: 0.3 to 1.5%,
Nb: 0.03-0.15%,
B: 0.0005 to 0.003%,
Al: 0.06% or less,
N: 0.006% or less,
A 490 MPa class high-strength steel excellent in high-temperature strength, characterized in that the balance consists of iron and inevitable impurities.
Ｎｉ：０．０５〜１．０％、
Ｃｕ：０．０５〜１．０％、
Ｃｒ：０．０５〜１．０％、
Ｔｉ：０．００５〜０．０２５％、
Ｖ：０．０１〜０．１％
の範囲で１種または２種以上を含有することを特徴とする請求項１に記載の高温強度に優れた４９０ＭＰａ級高張力鋼。In mass%,
Ni: 0.05 to 1.0%,
Cu: 0.05 to 1.0%,
Cr: 0.05 to 1.0%,
Ti: 0.005 to 0.025%,
V: 0.01 to 0.1%
The 490 MPa class high-strength steel excellent in high-temperature strength according to claim 1, comprising one or more kinds in the range of
Ｃａ：０．０００５〜０．００４％、
ＲＥＭ：０．０００５〜０．００４％、
Ｍｇ：０．０００１〜０．００６％
のいずれか１種または２種以上を含有することを特徴とする請求項１または２記載の高温強度に優れた４９０ＭＰａ級高張力鋼。In mass%,
Ca: 0.0005 to 0.004%,
REM: 0.0005 to 0.004%,
Mg: 0.0001 to 0.006%
The high-strength steel of 490 MPa class excellent in high-temperature strength according to claim 1 or 2, characterized in that any one or more of these are contained.
Ｍｏ／１５＋Ｖ／１０＋５Ｂ
と定義する溶接割れ感受性組成ＰCMが０．１８％以下であることを特徴とする高温強度及び溶接性に優れた４９０ＭＰａ級高張力鋼。Made of steel component according to any one of claims 1~3, P CM = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 +
Mo / 15 + V / 10 + 5B
490MPa grade high-tensile steel weld crack susceptibility composition P CM defining has excellent high-temperature strength and weldability and equal to or less than 0.18% and.
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| JP2003086307A JP4348102B2 (en) | 2002-05-20 | 2003-03-26 | 490 MPa class high strength steel excellent in high temperature strength and manufacturing method thereof |
| CN 03807395 CN1643167A (en) | 2002-03-29 | 2003-03-28 | High tensile steel excellent in high temperature strength and method for production thereof |
| EP03715623A EP1493828A4 (en) | 2002-03-29 | 2003-03-28 | High tensile steel excellent in high temperature strength and method for production thereof |
| PCT/JP2003/004040 WO2003087414A1 (en) | 2002-03-29 | 2003-03-28 | High tensile steel excellent in high temperature strength and method for production thereof |
| KR1020047015211A KR100630402B1 (en) | 2002-03-29 | 2003-03-28 | High tensile steel excellent in high temperature strength and method for production thereof |
| TW92107293A TWI235769B (en) | 2002-03-29 | 2003-03-31 | A high tensile strength steel excellent in high temperature and a method for producing the same |
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| JP2003086307A JP4348102B2 (en) | 2002-05-20 | 2003-03-26 | 490 MPa class high strength steel excellent in high temperature strength and manufacturing method thereof |
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| JP4466196B2 (en) * | 2004-05-24 | 2010-05-26 | 住友金属工業株式会社 | Steel sheet with excellent fatigue crack growth resistance and method for producing the same |
| JP4864297B2 (en) * | 2004-07-21 | 2012-02-01 | 新日本製鐵株式会社 | 490 MPa class high strength steel for welded structure excellent in high temperature strength and method for producing the same |
| JP4571915B2 (en) * | 2006-02-08 | 2010-10-27 | 新日本製鐵株式会社 | Refractory thick steel plate and manufacturing method thereof |
| KR101220618B1 (en) | 2010-12-27 | 2013-01-10 | 주식회사 포스코 | Flux cored arc weld wire having excellent workability and low-temperature toughness in weld metal joint and weld metal joint using the same |
| CN115094339B (en) * | 2022-07-27 | 2023-09-08 | 马鞍山钢铁股份有限公司 | Hot-rolled dual-phase high corrosion-resistant steel plate with tensile strength of 900MPa and manufacturing method thereof |
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