JP3922805B2 - Manufacturing method of high-tensile steel with excellent low-temperature toughness - Google Patents
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Description
【０００１】
【発明の属する技術分野】
本発明は、低温靭性が必要とされる構造部材に用いられる高張力鋼材の製造方法に関するものである。この方法で製造した鋼材は、例えば、海洋構造物、圧力容器、造船、橋梁、建築物、ラインパイプなどの溶接鋼構造物一般に用いることができるが、特に低温靭性を必要とする海洋構造物、造船等の構造物用鋼板として有用である。また、その他、構造部材として用いられ、低温靭性が要求される鋼管素材、あるいは形鋼にも適用可能である。
【０００２】
【従来の技術】
厚鋼板に代表される高張力鋼材は構造部材として用いられるため、構造物の安全性確保の観点から、低温靭性を要求される。高張力鋼材において、低温靭性を向上させる方法は種々提案されているが、他の特性劣化を生じることなく低温靭性を向上させる方法としては、結晶粒径の微細化がほぼ唯一の方法である。
【０００３】
最近、本発明者らはすでにフェライト（αとも称す）粒径を３μｍ程度以下に超細粒化して低温靱性を飛躍的に向上させる技術を、特開平８−６０２３９号公報、特開平１０−１２１１３２号公報等で開示している。本技術の要件はいずれも、オーステナイト（γとも称す）域及びγ／α二相域における加熱・熱間圧延条件の最適化により、変態αを微細化した上で、該αを加工により再結晶せしめて超微細粒を得ることにある。
該技術は、実際上は１〜３μｍ程度の超細粒化を達成するために適した手段であり、２ｍｍＶノッチシャルピー衝撃試験の破面遷移温度(vＴrS）で−１００℃〜−１２０℃程度の靱性を安定的に得るには十分な技術であるが、さらにそれ以上の低温靱性を達成するためには、一層の結晶粒径の微細化が必要となる。
【０００４】
現在、本発明者らがすでに開示している、αへの強加工とその結果としてのαの加工・再結晶と類似の方法による超細粒化技術により１μｍオーダーの超細粒化が達成されつつある。例えば、日本鉄鋼協会第１３５回春季講演大会論文集（ＣＡＭＰ−ＩＳＩＪ，ｖｏｌ．１１，Ｎｏ．３，１９８８）５６６頁に示されている、変形方向を変化させた他パス温間加工による方法で１μｍ程度のα粒径が達成されている。しかし、いずれも、実験室的な検討に止まっており、実用的、工業的な手段としては必ずしも確立していない。
【０００５】
【発明が解決しようとする課題】
本発明は、本発明者らがすでに開示している加工・再結晶による超細粒化技術をさらに発展させることにより、特殊な設備や、複雑な熱間加工あるいは熱処理工程を必要とせずに、かつ、特殊なせまい範囲の化学組成に限定することなく、平均α粒径が１μｍ以下の超細粒α組織を鋼材の全厚にわたって達成し、 vＴrSが−１５０℃以下の極めて優れた靱性を有する、引張強度が約５００MPa から９５０MPa までの高張力鋼材を製造するための方法を提供するものである。
【０００６】
【課題を解決するための手段】
今までの本発明者らの研究結果によれば、熱間圧延段階でのαの加工・再結晶によって得られる結晶粒径に支配的な影響を及ぼすのは、γ／α二相域〜α域での累積圧下率と加工前の組織の微細さである。従来は、αの超細粒化である限りは、αへの加工前に変態によってαを十分生成させておく必要があると考えられていた。即ち、比較的焼入性の低い鋼においてか、あるいは冷却を制御することにより塊状のαを微細に生成させ、該α主体組織に累積圧下率の大きい加工を加えることで超細粒化を達成した。ただし、該方法によっては、通常得られるα粒径は１μｍ〜３μｍの範囲で、非常に極端な熱履歴や過酷な加工条件によれば、１μｍ以下の結晶粒径も得られるが、工業的に安定して１μｍ以下を達成することは容易でないことも判明した。
【０００７】
そこで、本発明者らはさらに詳細な研究を行った結果、結晶粒径に対しては、前組織の微細化や熱間加工の影響が大きいのは当然のことながら、さらに、前組織の種類も大きな影響を及ぼし、１μｍ以下の超細粒組織を安定的に得るためには、微細な塊状α主体組織よりも、低温で制御圧延した転位を多量に含む未再結晶γから変態した微細ベイナイトを主体とする組織に加工を加えて再結晶させる方が有利であることを見出した。
【０００８】
さらに上記知見に加えて、細粒の未再結晶γからでも塊状αでなく、ベイナイト変態を確実に生じさせるには鋼の焼入性を一定以上高める必要があること、αとは異なった構造を有するベイナイトを再結晶・超細粒化するには、前組織が塊状αの場合とは異なった加工条件を設定する必要があること、等も判明し、１μｍ以下の超細粒組織を得るための新たな発明に至った。その要旨は以下に示す通りである。
【０００９】
（１）質量％で、
Ｃ ：０．０１〜０．２５％
Ｓｉ：０．０３〜１．０％
Ｍｎ：０．３０〜３．０％
Ａｌ：０．００３〜０．１％
Ｎ ：０．００１〜０．０１％
不純物として、
Ｐ ：０．０１％以下
Ｓ ：０．００５％以下、
を含有し、さらに
Ｃｒ：０．０１〜２．０％
Ｎｉ：０．０１〜６．０％
Ｍｏ：０．０１〜２．０％
Ｃｕ：０．０１〜１．５％
のうち、１種または２種以上を含有し、残部Ｆｅ及び不可避不純物からなり、かつ、下記（式１）式で示される理想焼入臨界直径（ＤＩ）が、２５．４〜７６２ｍｍである鋼片の熱間圧延にあたって、
Ａｃ３変態点以上、１１５０℃以下の温度に加熱した後、
累積圧下率が５５〜８０％で、圧延開始温度が８５０℃以下、終了温度が７５０℃以上の第１の圧延を行った後、
１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の第２の圧延を行うことを特徴とする、低温靭性に優れた高張力鋼材の製造方法。
Ｄ Ｉ （ｍｍ）＝１２．７・（Ｃ％）０．５・（１＋０．６４・Ｓｉ％）
・（１＋４．１０・Ｍｎ％）・（１＋０．２７・Ｃｕ％）
・（１＋０．５２・Ｎｉ％）・（１＋２．３３・Ｃｒ％）
・（１＋３．１４・Ｍｏ％）・（１＋１．６０・Ｗ％）
・・・ （式１）
【００１０】
（２）前記第１の圧延の前処理として、鋼片をＡｃ３変態点以上、１０５０℃以下の温度に加熱し、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却することを特徴とする、前記（１）に記載の低温靭性に優れた高張力鋼材の製造方法。
【００１１】
（３）前記第１の圧延の前処理として、鋼片を、Ａｃ３変態点以上、１１５０℃以下の温度に加熱した後、累積圧下率が２０〜５０％の予備熱間圧延を行い、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却することを特徴とする、前記（１）に記載の低温靭性に優れた高張力鋼材の製造方法。
【００１２】
（４）質量％で、
Ｃ ：０．０１〜０．２５％
Ｓｉ：０．０３〜１．０％
Ｍｎ：０．３０〜３．０％
Ａｌ：０．００３〜０．１％
Ｎ ：０．００１〜０．０１％
不純物として、
Ｐ ：０．０１％以下
Ｓ ：０．００５％以下、
を含有し、さらに
Ｃｒ：０．０１〜２．０％
Ｎｉ：０．０１〜６．０％
Ｍｏ：０．０１〜２．０％
Ｃｕ：０．０１〜１．５％
のうち、１種または２種以上を含有し、残部Ｆｅ及び不可避不純物からなり、かつ、下記（式１）式で示される理想焼入臨界直径（Ｄ Ｉ ）が、２５．４〜７６２ｍｍである鋼片の熱間圧延に先立ち、
該鋼片を、Ａｃ ３ 変態点以上、１０５０℃以下の温度に加熱し、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する前処理を施した後、
オーステナイト分率が５０〜９０％の二相域に加熱した後に、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の圧延を行うことを特徴とする、低温靭性に優れた高張力鋼材の製造方法。
Ｄ Ｉ （ｍｍ）＝１２．７・（Ｃ％）０．５・（１＋０．６４・Ｓｉ％）
・（１＋４．１０・Ｍｎ％）・（１＋０．２７・Ｃｕ％）
・（１＋０．５２・Ｎｉ％）・（１＋２．３３・Ｃｒ％）
・（１＋３．１４・Ｍｏ％）・（１＋１．６０・Ｗ％）
・・・ （式１）
（５）前記前処理に代えて、前記鋼片を、Ａｃ ３ 変態点以上、１１５０℃以下の温度に加熱した後、累積圧下率が２０〜５０％の予備熱間圧延を行い、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する前処理を施すことを特徴とする、前記（４）に記載の低温靭性に優れた高張力鋼材の製造方法。
【００１３】
（６）最後の熱間圧延終了後の鋼材を２〜４０℃／ｓの冷却速度で２０℃〜５００℃まで加速冷却することを特徴とする、前記（１）〜（５）のいずれかに記載の低温靭性に優れた高張力鋼材の製造方法。
【００１４】
（７）４５０℃以上、Ａｃ１変態点以下で焼戻しを行うことを特徴とする、前記（６）に記載の低温靭性に優れた高張力鋼材の製造方法。
【００１５】
（８）質量％で、
Ｔｉ：０．００３〜０．１０％
Ｖ ：０．００５〜０．５０％
Ｎｂ：０．００３〜０．１０％
Ｚｒ：０．００３〜０．１０％
Ｔａ：０．００５〜０．２０％
Ｗ ：０．０１〜２．０％
Ｂ ：０．０００３〜０．００２０％
の１種または２種以上をさらに含有することを特徴とする、前記（１）〜（７）のいずれかに記載の低温靭性に優れた高張力鋼材の製造方法。
【００１６】
（９）質量％で、
Ｍｇ：０．０００５〜０．０１％
Ｃａ：０．０００５〜０．０１％
ＲＥＭ：０．００５〜０．１０％
のうち１種または２種以上をさらに含有することを特徴とする、前記（１）〜（８）のいずれかに記載の低温靭性に優れた高張力鋼材の製造方法。
【００１７】
【発明の実施の形態】
以下に本発明について詳細に説明する。
本発明は加工・再結晶によるαの超細粒化方法において、従来達成レベルを凌駕するための新しい手段として、ベイナイト主体組織に加工を加えることに特徴を有する。ベイナイト組織はラス構造を有するため、再結晶のしやすさ、再結晶後の粒径を支配する観点での実質的な組織の微細さという意味では塊状α組織に比べて格段に有利である。ただし、最終的なα粒径を１μｍ以下とするためには、単にベイナイト組織としただけでは十分でなく、粒径が微細で、かつ転位密度の高い未再結晶γから変態させたベイナイトとする必要があることを詳細な実験により見出した。
【００１８】
細粒未再結晶γから安定的にベイナイト変態させるためには、鋼の組成も適正化する必要があり、強度・靱性確保のための個々の元素の限定とともに、焼入性の指標である理想焼入臨界直径（ＤI ）を一定以上にすることも重要である。即ち、化学組成と製造条件との最適な組み合わせが必要となる。以下、本発明の詳細な説明を製造条件と化学組成とに分けて行う。
【００１９】
先ず、本発明の製造条件について説明する。
本発明の方法は、先ず、「個々の元素が適正な化学組成を有し、かつ（式１）式で示すＤＩ 値が適正範囲の鋼片をＡｃ３変態点以上、１１５０℃以下の温度に加熱した後、累積圧下率が５５〜８０％で、圧延開始温度が８５０℃以下、終了温度が７５０℃以上の第１の圧延を行った後、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の第２の圧延を行うこと」を第１の方法とする。さらに、一層の超細粒化と、均一性の確保のために、鋼片に対して、「鋼片をＡｃ３変態点以上、１０５０℃以下の温度に加熱し、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する」か、もしくは「鋼片を、Ａｃ３変態点以上、１１５０℃以下の温度に加熱した後、累積圧下率が２０〜５０％の圧延を行い、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する」前組織微細化処理を行った後に第１の方法を行う方法を第２の方法とする。また、さらに、該前組織微細化処理を行うことを前提として、「鋼片をオーステナイト分率が５０〜９０％の二相域に再加熱し、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の圧延を行う」方法を第３の方法とする。
【００２０】
なお、本発明における鋼片とは、溶鋼を鋳型で凝固させたインゴットや連続鋳造法により製造したスラブ、ビレット、ブルーム等の鋳片及びこれらの鋳片を、形状調整や偏析軽減を主目的として、分塊圧延、熱間圧延した鋼片等、本発明の要件となる熱間圧延、熱処理を施される前の鋳片、鋼片、鋼板全般を指す。
【００２１】
以上に記載の３つの方法が本発明の基本要件であるが、強度、靱性等の材質調整の必要に応じて、該熱間圧延の終了後、「２〜４０℃／ｓの冷却速度で２０℃〜３００℃まで加速冷却」したり、「４５０℃以上、Ａｃ1 変態点以下」で焼戻しを行うことも可能である。
【００２２】
上記、本発明の基本要件である３つの方法について、さらに詳細に説明する。
先ず、第１の方法について説明する。鋼片を熱間圧延に先だって加熱するが、該加熱温度はＡｃ3 変態点以上、１１５０℃以下の範囲とする必要がある。これは、加熱温度がＡｃ3 変態点未満であると、γ化されない部分が粗大α組織となって、その後の熱間圧延によっても超細粒化されず、そのために最終組織が顕著に混粒となるためであり、逆に１１５０℃超であると、加熱γ粒径が粗大となって、変態前のγの微細化が不十分となり、最終的な超細粒化も不十分となるためである。
【００２３】
鋼片をＡｃ3 変態点以上、１１５０℃以下で加熱した後、累積圧下率が５５〜８０％で、圧延開始温度が８５０℃以下、終了温度が７５０℃以上の第１の圧延を行って変態前のγの粒径、形態を調整する。即ち、γの細粒化と加工歪の十分な導入のためには累積圧下率は５５％以上必要である。累積圧下率は大きいほど細粒化や転位の導入に有利となるが、該圧延での累積圧下率が過大であると、引き続き行う第２の圧延での圧下率の確保が困難となるため、圧延効果が飽和する８０％を第１の圧延における上限とする。一方、累積圧下率が５５％未満であると、未再結晶γへの歪の蓄積が不十分で、変態後のベイナイト主体組織の細分化が不十分となり、最終的な結晶粒径の微細化も不十分となる。また、第１の圧延温度は開始温度を８５０℃以下、終了温度を７５０℃と限定するが、これは、８５０℃超の圧延はγの細粒化や転位の導入に対する効果が小さいため、圧延の効果を十分発揮させるためには圧延開始温度を８５０℃以下とすることが有利であるためであり、終了温度が７５０℃未満であると、圧延の効果が飽和する一方で、第２の圧延温度確保が困難となるためである。
【００２４】
圧延温度が本発明範囲の８５０℃〜７５０℃であれば、各パスの圧延効果はほぼ累積されるため、各パスの圧下率は限定する必要はなく、第１の圧延の効果は累積圧下率が本発明の範囲を満足すれば達成される。
なお、本発明においては、該圧延の前に８５０℃超の温度で形状調整等を主目的とした圧延を加えることを妨げるものではない。即ち、第１及び第２の圧延での累積圧下率が確保でき、本発明の効果を損なわない範囲であれば、８５０℃超での加工を加えても構わない。
【００２５】
上記第１の圧延によって変態前のγの粒径、形態を調整して、変態後のベイナイト主体組織の微細化を図った後、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の第２の圧延を行う。
第２の圧延によってベイナイト主体組織に加工が加えられ、転位の導入・回復・再結晶過程を通って、αだけでなく、ベイナイト組織も細分化され、全体がαの超細粒組織となる。そのためには、先ず、γを十分変態させて第２の圧延に入る前に少なくとも３０％以上ベイナイト組織としておく必要がある。残部はα、あるいはγその他の組織でも構わない。第１の圧延を施しておけば、ベイナイトだけでなくその他の変態組織も十分微細化されるため、超細粒化に問題はない。また、未変態γが存在していても圧延中にベイナイトへ変態するため、これも問題はない。ただし、該組織形態を確保するためには、圧延開始温度を７００℃以下とする必要がある。７００℃超では、本発明の化学組成全般にわたって安定して上記加工前組織形態を保証できない。また、圧延終了温度は６００℃以上とする必要がある。これは、６００℃未満であると、ベイナイト組織の再結晶が困難となって超細粒化が十分でない場合が多くなるためと、変形抵抗が過大となって圧延機への負荷が過大となるためと、鋼材の形状確保が困難となるためである。
【００２６】
上記の理由により第２の圧延の温度は７００℃〜６００℃の範囲とする必要があるが、累積圧下率及び各パスの圧下量も規定する必要がある。即ち、ベイナイト主体組織を７００℃〜６００℃での圧延において十分再結晶させて超細粒化させるためには、累積圧下率を５０〜９５％とすることが先ず前提となる。即ち、累積圧下率が５０％未満であると、各パスの圧下条件によらず再結晶が困難で、また再結晶したとしても達成される結晶粒径に限度があるため、累積圧下率は５０％以上が必要である。累積圧下率は大きいほど細粒化には有利であるが、９５％超では細粒化効果が飽和する一方で、温度の確保や生産性に問題を生ずることから、本発明では上限を９５％とした。
【００２７】
第２の圧延において上記の理由により、圧延温度と累積圧下率を限定するが、ベイナイト主体組織に加工を加えてα粒径を１μｍ以下にするためには、圧延温度と累積圧下率の限定だけでは十分でなく、累積圧下率が５０から９５％の１パスあるいは多パス圧延を行うに際して、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含む必要がある。これは、本発明ではベイナイト主体組織に加工を加える必要上、圧延温度がかなり低くなるが、その場合にはα組織主体に加工を加える場合よりも再結晶が困難となるため、圧延中のいずれかのパスで大圧下の加工を加えることで再結晶の促進を図る必要がある。また、結晶方位によって再結晶の難易度が異なり、方位によっては累積圧下率を高めただけでは最後まで再結晶しないため、混粒組織となりやすいが、このような混粒組織を解消するためにも圧延のいずれかの段階で大圧下圧延を行うことが好ましい。そのための下限の圧下率は３０％である。この場合、加工率が大きいこと自体と加工発熱により温度が上昇することで組織の回復・再結晶が促進され、組織の均一微細化が図られる。
【００２８】
該圧下率は大きければ大きい程、細粒化には好ましいが、
（１）過剰な加工発熱により再結晶したα粒が粒成長して粗大化する。
（２）１パスの圧下率が過大であると圧延機に負荷がかかりすぎる。
（３）鋼板形状に悪影響を及ぼす。
という３つの理由から、１パスあたりの圧下率の上限は７０％とする。１パスの圧下率が３０〜７０％の圧延はα粒径の微細化の目的から、圧延の初期の方が効果が大であり、また好ましいが、累積圧下率、圧延温度域が本発明範囲であれば、圧延のどの段階で行っても必要な超細粒組織は得られる。また、３０〜７０％圧下のパスも多い方が好ましいが、最低１パス含まれていれば超細粒化に支障はない。
【００２９】
以上が、「個々の元素が適正な化学組成を有し、かつ（式１）式で示すＤＩ値が適正範囲の鋼片をＡｃ３変態点以上、１１５０℃以下の温度に加熱した後、累積圧下率が５５〜８０％で、圧延開始温度が８５０℃以下、終了温度が７５０℃以上の第１の圧延を行った後、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の第２の圧延を行うこと」を特徴とする第１の方法の限定理由である。次に、第３の方法の限定理由を前組織微細化処理の理由とともに示す。なお、前組織微細化処理は第１の方法においても有効で、第１の方法での効果を妨げるものでなく、第１の方法による圧延の前に前組織微細化処理を施すことで、第１の方法以上の超細粒化、混粒度の低減により、第１の方法以上の靱性を得るために有効な方法が第２の方法である。
【００３０】
第３の方法は、「請求項２または３に記載の前処理と同じ前処理により前組織を微細化した後、第１の方法のような第１の圧延を行うことなく、オーステナイト分率が５０〜８０％の二相域に再加熱し、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の圧延を行う」ことを特徴とするものである。即ち、第１の方法と異なり、熱間圧延の前にγ域とα域の中間の二相域に加熱するが、これは、前組織を請求項２または３に記載の前処理と同じ前処理で微細化した場合には、本発明のように、微細・未再結晶γでベイナイト主体組織となる焼入性の鋼においては、直接二相域に再加熱した方が、圧延中に微細・未再結晶γからベイナイト変態が促進され、安定的に所望の組織とすることが可能であり、かつ、加熱温度が低い分、温度待ちの時間を短くでき、生産性の向上につながるためである。
【００３１】
前組織を微細化した上で二相域に再加熱する場合、本発明ではベイナイト主体組織とする必要性から、γ分率を一定範囲に制御する必要がある。即ち、γ分率が小さいと、α主体組織となるため、１μｍ以下の超細粒組織を得ることは困難となる。最終組織のα粒径を１μｍ以下とするために必要なベイナイト主体組織とするためのγ分率の下限は実験に基づいて、５０％と定める。一方、γ分率が多すぎると、実質的にγ単相と変わらなくなるため、ベイナイト変態促進の効果を有するγ分率が９０％を上限とする。
【００３２】
以上のように、請求項２または３に記載の前処理と同じ前処理により前組織を微細化した後についての圧延条件は第１の方法の第２の圧延と全く同じである。
次に、第１の方法においては一層の超細粒化を図るために必要に応じて採用される工程であるが、第３の方法においては超細粒化に必須の工程となる、前組織微細化のための前処理について説明する。
【００３３】
請求項２または４に記載の前処理は、鋼片の粗大な凝固組織を解消して、超細粒化のための二相域〜α域での加工に入る前の組織を微細化する。その際、鋼片をＡｃ３変態点以上、１０５０℃以下の温度に加熱するが、これは、加熱温度がＡｃ３変態点未満であるとその後の冷却条件如何によらず粗大な組織が残存し、１０５０℃超であると加熱時のγ粒径が粗大化して冷却後の組織の微細化が不十分となるためである。また、加熱保持後の冷却も０．２〜２０℃／ｓの冷却速度で３００℃以下まで制御する必要がある。該条件は加熱温度を前記のように規定して加熱γ粒径を微細化しても、その後の冷却による変態組織を微細化するために必要である。
【００３４】
即ち、鋼片は一般的に厚いために、放冷程度でも鋼片内部の冷却は徐冷となる場合が多く、その場合には変態組織は粗大となる。そのため、冷却後の組織を微細化するためには０．２℃／ｓ以上で冷却する必要がある。冷却速度は大きければ大きいほど組織微細化には好ましいが、板厚の大きい鋼片では冷却速度を極端に大きくすることは現実的でなく、２０℃／ｓ以上であれば最終的な加工・再結晶組織に有害な極端な粗大組織の出現は抑制できることから、本発明では上限を２０℃／ｓとした。また、このような粗大組織の出現は、該制御冷却を３００℃まで実施すればその後の冷却条件によらずに抑制できる。なお、鋼片厚さが比較的小さく、空冷によっても０．２℃／ｓ以上で冷却される場合には当然空冷でも構わない。
【００３５】
上記、請求項２または４に記載の前処理では熱処理のみにより組織微細化を図るが、請求項３または５に記載の前処理では熱間圧延により鋼片組織を微細化する。熱処理のみならず圧延による加工をも伴う方が組織微細化には有利であるが、鋼片厚みを減ずることにもなるため、鋼片厚みが仕上げ板厚に対して厚く、累積圧下率の裕度が高い場合に好ましい方法と言える。
【００３６】
該方法において、先ず鋼片をＡｃ３変態点以上、１１５０℃以下の温度に再加熱する。Ａｃ３変態点以上に再加熱するのは請求項２または４に記載の前処理の場合と同じ理由であり、一方、後の圧延がγの細粒化効果を有するため、再加熱温度の上限は圧延のない場合に比べて緩和されるが、再加熱γ粒径が過大であると、後の圧延によっても細粒化が不十分となるため、本発明では上限を１１５０℃とする。
【００３７】
Ａｃ３変態点以上、１１５０℃以下の温度に再加熱した後、γ組織微細化のために熱間圧延を行うが、その際、該熱間圧延の累積圧下率を２０〜５０％とする。これはＡｃ３変態点〜１１５０℃の再加熱に続く圧延の場合に、累積圧下率が２０％未満であると、折角圧延を行ってもγ粒の微細化とその結果としての変態組織の微細化が不十分であり、一方、５０％超では、最終的な超細粒化のための圧延段階での累積圧下率が十分とれなくなるためである。熱間圧延後の冷却は請求項２または４に記載の前処理と同じで、変態組織を十分微細化するための条件として、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する。
【００３８】
以上が、本発明における基本要件についての説明であるが、その他、製造方法に関する要件の限定理由を以下に付け加える。
即ち、鋼材の強度調整、靱性向上を目的として、最終の圧延終了後の鋼材を２〜４０℃／ｓの冷却速度で２０℃〜５００℃まで加速冷却することができる。加速冷却により内部の組織が微細化するとともに、第二相がより硬質なものへと変化することによる。この効果を発揮するためには冷却速度は２℃／ｓ以上必要である。冷却速度は大きい方が組織微細化、硬質相形成には有利であるが、厚手材では無制限に冷却速度を大きくすることは実用上困難であることから、組織制御に対する加速冷却効果が飽和しない範囲として本発明においてはその上限を４０℃／ｓとする。また、該加速冷却は２０℃〜５００℃まで行うことが好ましい。即ち、２０℃未満まで冷却しても組織制御に有効でなく、加速冷却の停止温度が５００℃超では、組織微細化が十分でなく、また硬質相が形成され難いために強度調整に有効でなく、靱性が劣化する場合があるためである。また、本発明で規定した加速冷却条件によれば、一旦形成された表層部の超細粒組織の粒成長抑制にも有効である。
【００３９】
またさらに、強度調整、靭性向上、形状改善の目的で、焼き戻し処理を施すことも可能である。その場合には、表層部に形成された超細粒組織を損なわないことが必須要件となる。本発明では焼き戻し温度を、４５０℃〜Ａｃ１変態点に限定するが、これは、４５０℃未満では焼き戻しの効果が明確ではなく、Ａｃ１変態点超では表層部の超細粒組織の形態を損なうためである。ただし、超細粒層の粒成長抑制をより確実にするためには、焼戻し温度は７００℃を超えないことがより好ましい。なお、本発明の焼き戻し温度範囲であれば、焼き戻しの加熱保持時間は任意であるが、同様に表層部の超細粒組織保存の観点からは、保持時間は５時間以内であることが好ましい。
【００４０】
以上が、本発明の低温靱性に優れた高張力鋼材の製造方法に関する要件であるが、該製造方法により効果を発揮するためには個々の化学成分についても下記に述べる理由により、各々限定する必要がある。
【００４１】
即ち、Ｃは、鋼の強度を向上させる有効な成分として含有するもので、０．０１％未満では構造用鋼に必要な強度の確保が困難であるが、０．２５％を超える過剰の含有は母材及び溶接部の靭性や耐溶接割れ性を低下させるので、０．０１〜０．２５％の範囲とした。
【００４２】
次に、Ｓｉは、脱酸元素として、また、母材の強度確保に有効な元素であるが、０．０３％未満の含有では脱酸が不十分となり、また強度確保に不利である。逆に１．０％を超える過剰の含有は粗大な酸化物を形成して延性や靭性の劣化を招く。そこで、Ｓｉの範囲は０．０３〜１．０％とした。
【００４３】
また、Ｍｎは、母材の強度および靭性の確保に必要な元素であり、最低限０．３％以上含有する必要があるが、過剰に含有すると、硬質相の生成や粒界脆化等により母材靱性や溶接部の靭性、さらに溶接割れ性など劣化させるため、材質上許容できる範囲で上限を３．０％とした。
【００４４】
Ａｌは、脱酸やオーステナイト粒径の細粒化等に有効な元素であるが、効果を発揮するためには０．００３％以上含有する必要がある。一方、０．１％を超えて過剰に含有すると、粗大な酸化物を形成して延性を極端に劣化させるため、０．００３％〜０．１％の範囲に限定する必要がある。
【００４５】
Ｎは、ＡｌやＴｉと結びついてγ粒微細化に有効に働くため、微量であれば機械的特性に有効に働く。また、工業的に鋼中のＮを完全に除去することは不可能であり、必要以上に低減することは製造工程に過大な負荷をかけるため好ましくない。そのため、工業的に制御が可能で、製造工程への負荷が許容できる範囲として下限を０．００１％とする。過剰に含有すると、固溶Ｎが増加し、延性や靭性に悪影響を及ぼす可能性があるため、許容できる範囲として上限を０．０１％とする。
【００４６】
本発明では、微細・未再結晶γから確実に微細なベイナイト組織を変態させるために、後述のように臨界焼入臨界直径（ＤI ）を適正にする必要があるが、他の特性に影響を与えずにＣ、Ｓｉ、Ｍｎだけで所望のＤI 値に調整することは困難であるので、焼入性への寄与の大きな、Ｃｒ、Ｎｉ、Ｍｏ、Ｃｕのうちの１種または２種以上の添加によりベイナイト組織化させることが必要である。ただし、個々の添加量についても、下記のように限定する必要がある。
【００４７】
Ｃｒ及びＭｏは、いずれも母材の強度向上に有効な元素であるが、明瞭な効果を生じるためには０．０１％以上必要であり、一方、２．０％を超えて添加すると、靭性及び溶接性が劣化する傾向を有するため、各々０．０１〜２．０％の範囲とする。
【００４８】
また、Ｎｉは、母材の強度と靭性を同時に向上でき、非常に有効な元素であるが、効果を発揮させるためには０．０１％以上含有させる必要がある。含有量が多くなると強度、靭性は向上するが６．０％を超えて添加しても効果が飽和する一方で、溶接性が劣化するため、上限を６．０％とする。また、６．０％を超える添加ではＤI 値を本発明範囲内としてもベイナイト変態が抑制されて超細粒化に好ましくない効果も顕在化する。
【００４９】
次に、Ｃｕも、ほぼＮｉと同様の効果を有するが、１．５％超では熱間加工性に問題を生じるため、０．０１〜１．５％の範囲に限定する。
【００５０】
ＰおよびＳは、不純物元素で、延性、靭性を劣化させる元素であり、極力低減することが好ましいが、材質劣化が大きくなく、許容できる量として、Ｐの上限を０．０１％、Ｓの上限を０．００５％に限定する。
【００５１】
以上が本発明の鋼材の基本成分の限定理由であるが、本発明においては、強度・靭性の調整のために、必要に応じて、Ｔｉ、Ｖ、Ｎｂ、Ｚｒ、Ｔａ、Ｗ、Ｂの１種または２種以上を含有することができる。
【００５２】
Ｔｉは、析出強化により母材強度向上に寄与するとともに、ＴｉＮの形成により加熱オーステナイト粒径微細化にも有効な元素であり、靭性向上にも有効な元素であるが、効果を発揮するためには０．００３％以上の含有が必要である。一方、０．１０％を超えると、粗大な析出物、介在物を形成して靭性や延性を劣化させるため、上限を０．１０％とする。
【００５３】
Ｖも、ＶＮを形成して強度向上に有効な元素であるが、過剰の含有では析出脆化により靭性が劣化する。従って、靭性の大きな劣化を招かずに、効果を発揮できる範囲として、０．００５〜０．５０％の範囲に限定する。
【００５４】
Ｎｂは、Ｎｂ（Ｃ，Ｎ）を形成することで強度・靭性の向上に有効な元素であるが、過剰の含有では析出脆化により靭性が劣化する。従って、靭性の劣化を招かずに、効果を発揮できる範囲として、０．００３〜０．１０％の範囲に限定する。
【００５５】
Ｚｒも、窒化物を形成する元素であり、Ｔｉと同様の効果を有するが、その効果を発揮するためには０．００３％以上の含有が必要である。一方、０．１０％を超えると、Ｔｉと同様、粗大な析出物、介在物を形成して靭性や延性を劣化させるため、０．００３〜０．１０％の範囲に限定する。
【００５６】
Ｔａも、強度・靭性の向上に有効な元素であるが、効果を発揮するためには０．００５％以上の含有が必要である。一方、０．２０％を超えると、析出脆化や粗大な析出物、介在物による靭性劣化を生じるため、上限を０．２０％とする。
【００５７】
Ｗは、固溶強化及び析出強化により母材強度の上昇に有効であるが、効果を発揮するためには０．０１％以上必要である。一方、２．０％を超えて過剰に含有すると、靭性劣化が顕著となるため、上限を２．０％とする。
【００５８】
Ｂは、微量で確実にＮと結びつくため、固溶Ｎ固定により靭性向上や、焼入性向上による強度・靭性向上に有効な元素であるが、効果を発揮するためには０．０００３％以上必要である。一方、０．００２０％を超えて過剰に含有するとＢＮが粗大となり、延性や靭性に悪影響を及ぼす。また溶接性も劣化させるため、上限を０．００２０％とする。
【００５９】
さらに、母材靱性の向上、延性の向上、継手靭性の向上のために、必要に応じて、Ｍｇ、Ｃａ、ＲＥＭの１種または２種以上を含有することができる。
Ｍｇ、Ｃａ、ＲＥＭは、いずれも硫化物の熱間圧延中の展伸を抑制して延性特性向上に有効である。酸化物を微細化させて一層の超細粒化に寄与して母材靱性向上に有利となる上、継手靭性の向上にも有効に働く。その効果を発揮するための下限の含有量は、Ｍｇ及びＣａは０．０００５％、ＲＥＭは０．００５％である。一方、過剰に含有すると、硫化物や酸化物の粗大化を生じ、延性、靭性の劣化を招くため、上限を各々、Ｍｇ、Ｃａは０．０１％、ＲＥＭは０．１０％とする。
【００６０】
以上が、個々の元素の限定理由であるが、微細・未再結晶γから確実に微細なベイナイト組織を変態させるために、個々の元素の限定だけでなく、焼入性を適正化するために、焼入性の指標である、下記（式１）式に示す臨界焼入臨界直径（ＤＩ）を２５．４〜７６２ｍｍとする必要がある。
Ｄ Ｉ （ｍｍ）＝１２．７・（Ｃ％）０．５・（１＋０．６４・Ｓｉ％）
・（１＋４．１０・Ｍｎ％）・（１＋０．２７・Ｃｕ％）
・（１＋０．５２・Ｎｉ％）・（１＋２．３３・Ｃｒ％）
・（１＋３．１４・Ｍｏ％）・（１＋１．６０・Ｗ％）
・・・ （式１）
【００６１】
（式１）式で示されるＤＩ値が２５．４ｍｍ未満であると、焼入性が不足するため、変態組織中のベイナイト組織の割合が十分でなくなり、１μｍ以下の超細粒組織とすることが困難である。一方、ＤＩ値が７６２ｍｍを超えて過剰となると、前組織微細化や二相域加熱等を行っても、γからベイナイト変態せずに、マルテンサイト変態するため、変態温度が４００程度以下と非常に低くなって、変態組織に加工を加えることが実質的に不可能となる。以上の理由により本発明では、（式１）式に示す臨界焼入臨界直径（ＤＩ）を２５．４〜７６２ｍｍに限定する。
【００６２】
【実施例】
以上が、本発明の要件についての説明であるが、さらに、実施例に基づいて本発明の効果を示す。
表１に示す化学組成の鋼片を用いて、表２、表３に示す製造条件で鋼板を製造した。製造した鋼板の、本発明の特徴となるフェライト組織の特徴（鋼板板厚中心部の平均粒径、混粒度、粒形態等）や機械的性質（強度、２ｍｍＶノッチシャルピー衝撃特性、ＥＳＳＯ特性）の測定結果も合わせて表２、表３に示す。なお、表２は、請求項１〜３に関連した本発明鋼と比較例とを示した結果であり、表３は請求項４、５に関連した本発明鋼と比較例とを示した結果である。
【００６３】
組織観察、粒径の測定は、本発明鋼及び比較例のうちのフェライト粒径が比較的微細な鋼については倍率１０００倍から５０００倍の走査型電子顕微鏡（ＳＥＭ）組織写真に基づいて実施し、粒径が粗い比較例については倍率５００倍の光学顕微鏡組織写真に基づいて実施した。引張特性は圧延方向に直角な方向（Ｃ方向）の板厚中心部から丸棒引張試験片を採取して実施した。２ｍｍＶノッチシャルピー衝撃試験も圧延方向に直角な方向（Ｃ方向）の板厚中心部から試験片板厚１０ｍｍの標準試験片を採取して行った。シャルピー衝撃特性は破面遷移温度（ｖＴｒＳ）で評価した。さらに、脆性き裂伝播停止特性（アレスト性）を温度勾配型のＥＳＳＯ試験で調査した。ＥＳＳＯ特性は、Ｃ方向から試験片を採取し、Ｋｃａが４００ｋｇｆ／ｍｍ１．５ （＝３９２０Ｎ／ｍｍ １．５ ）となる温度（ＴＫｃａ ４００）で評価した。
【００６４】
表２のうちの試験Ｎｏ．Ａ１〜Ａ９と表３のうちの試験Ｎｏ．Ａ１０〜Ａ１６は、本発明の化学組成を有する鋼番１〜１１を用いて、本発明の製造方法により製造した鋼板であり、いずれも、平均フェライト粒径は１μｍ以下と、従来のフェライトの加工・再結晶による超細粒鋼のレベルよりも１段階さらに微細な超細粒組織が得られており、極めて良好な靱性、アレスト性が得られている。すなわち、従来の平均フェライト粒径が３μｍ程度以下の超細粒鋼の vＴrSが−１００〜−１５０℃程度であるのに対して、本発明の平均粒径が１μｍ以下の超細粒鋼の vＴrSは全て−１７０℃以下で、大部分は液体窒素温度（−１９６℃）でも延性破壊主体の破壊形態を示しており、また、このような極めて高い靱性を反映して、全厚でのＥＳＳＯ特性も、ＴKca ４００で−１３２℃以下と、極めて良好であり、鋼材の安全性が飛躍的に増加していることが明白である。
【００６５】
一方、表２、表３の結果から、本発明の範囲を逸脱している試験Ｎｏ．Ｂ１〜Ｂ１１の鋼板は本発明により製造された試験Ｎｏ．Ａ１〜Ａ１６の鋼板に比べて、靱性、アレスト性が大幅に劣っていることが明らかである。
試験Ｎｏ．Ｂ１〜Ｂ５、及びＮｏ．Ｂ１１は、化学組成が本発明を満足していないために、製造方法は本発明を満足しているものの、十分な靱性、アレスト性が達成できなかった例である。
【００６６】
すなわち、試験Ｎｏ．Ｂ１は、（式１）式で示すＤＩ値が過小で、二相域加工前の組織がベイナイト主体組織とならないために、最終的に得られるフェライト粒径が１．９８μｍと、１μｍ以下となっておらず、本発明鋼に比べてｖＴｒＳとアレスト性（ＴＫｃａ ４００）が劣る。
【００６７】
試験Ｎｏ．Ｂ２は、逆にＤI 値が過大で、焼入性が過大なために、請求項１に基づく方法で、第１の圧延を行い、さらに低温で圧延を加えているにもかかわらず、大部分が未変態オーステナイトのために伸長オーステナイトから変態したベイナイト主体組織となっており、そのために超細粒フェライト組織が得られておらず、従って、本発明鋼に比べて vＴrSとアレスト性（ＴKca ４００）が大幅に劣る。
【００６８】
また、試験Ｎｏ．Ｂ１１は試験Ｎｏ．Ｂ２と同じ鋼片を用いて、請求項５に基づいた方法で製造したものであり、二相域加熱を施しているが、二相域加熱時にオーステナイト相であった部分は、加工前、加工中にも変態を生ぜず、最終組織は伸長したベイナイト主体組織となって、Ｎｏ．Ｂ２と同様、超細粒化は達成できておらず、靱性、アレスト性も本発明鋼に比べて劣る。
【００６９】
試験Ｎｏ．Ｂ３は、Ｃが過剰なため、超細粒化は十分であるにもかかわらず、良好な靱性レベルが達成されない。
試験Ｎｏ．Ｂ４は、Ｐ、Ｓがともに過剰なため、超細粒化は十分であるにもかかわらず、本発明鋼に比べて vＴrS、ＴKca ４００ともに劣る。
試験Ｎｏ．Ｂ５は、Ｃｒ、Ｍｏがともに過剰なため、超細粒化は十分であるにもかかわらず、良好な靱性レベルが達成されない。
【００７０】
一方、試験Ｎｏ．Ｂ６、Ｂ７、Ｂ８〜Ｂ１０は、化学組成は本発明を満足しているが、製造法が本発明の範囲を逸脱しているために、本発明の目的としている、１μｍ以下の平均フェライト粒径が達成できておらず、本発明により製造した鋼板に比べて特性が劣っている例である。
【００７１】
試験Ｎｏ．Ｂ６、Ｂ７は、請求項１に関連した比較例であり、Ｎｏ．Ｂ６は、第２の圧延において、いずれの圧延パスでも圧下率が３０％以上となっていないために、また、Ｎｏ．Ｂ７は、第１の圧延の累積圧下率が本発明の範囲を満足していないために、いずれも最終的なフェライト粒径が本発明による鋼板に比べて粗大で、靱性、アレスト性も本発明鋼に比べて劣る。
【００７２】
試験Ｂ８〜Ｂ１０は、請求項４、５に関連した比較例である。
試験Ｎｏ．Ｂ８は、二相域加熱・圧延前の前組織微細化工程が施されていないために、二相域加熱段階で粗大な組織が残存し、最終的な組織も混粒・粗大で、靱性、アレスト性が劣る。
試験Ｎｏ．Ｂ９は、二相域加熱後の圧延において、いずれの圧延パスでも圧下率が３０％以上となっていないために、本発明鋼に比べて超細粒化が十分でなく、ｖＴｒＳ、ＴＫｃａ ４００とも従来レベルに止まっている。
【００７３】
試験Ｎｏ．Ｂ１０は、二相域加熱温度が低いために、加熱段階でのオーステナイト分率が少なく、その結果、１μｍ以下のフェライト粒径を達成するために必須の要件である、「転位を多量に含む未再結晶γから変態した微細ベイナイトを主体とする組織に加工を加える」条件を満足できていないため、平均フェライト粒径が２μｍ程度までしか細粒化できず、従来のフェライトの加工・再結晶による超細粒鋼の靱性レベルしか得られていない。
【００７４】
以上の実施例から、本発明によれば、平均フェライト粒径が１μｍ以下の超細粒フェライト組織を鋼材の全厚にわたって達成でき、極めて良好なシャルピー衝撃特性及びアレスト性を有する高張力鋼材の製造が可能になることは明白である。
【００７５】
【表１】
【表２】
【表３】
【表４】
【表５】
【表６】
【表７】
【表８】
【発明の効果】
本発明は、平均フェライト粒径が１μｍ以下の超細粒フェライト組織を鋼材の全厚にわたって達成させて、 vＴrSが−１５０℃以下の極めて優れた低温靱性と脆性き裂伝播停止特性とを有する、引張強度が約５００MPa から９５０MPa 超級までの高張力鋼材を製造するための方法を提供するものであり、構造物の安全性を飛躍的に高めることが可能な高張力鋼材を経済性、生産性を損なうことなく製造できる手段として、産業上の価値の極めて高い発明であると言える。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for producing a high-tensile steel material used for a structural member requiring low temperature toughness. Steel materials produced by this method can be used for general welded steel structures such as marine structures, pressure vessels, shipbuilding, bridges, buildings, line pipes, etc., especially marine structures that require low temperature toughness, It is useful as a steel plate for structures such as shipbuilding. In addition, the present invention can also be applied to steel pipe materials or shape steels that are used as structural members and require low temperature toughness.
[0002]
[Prior art]
Since high-tensile steel materials typified by thick steel plates are used as structural members, low temperature toughness is required from the viewpoint of ensuring the safety of structures. Various methods for improving the low temperature toughness in high-tensile steel materials have been proposed, but the refinement of crystal grain size is almost the only method for improving the low temperature toughness without causing other characteristic deterioration.
[0003]
Recently, the present inventors have already developed a technique for dramatically improving the low-temperature toughness by reducing the ferrite (also referred to as α) grain size to about 3 μm or less and improving the low-temperature toughness. It is disclosed in No. Gazette. The requirements of this technology are that the transformation α is refined by optimizing the heating and hot rolling conditions in the austenite (also referred to as γ) region and the γ / α two-phase region, and the α is recrystallized by processing. It is to obtain ultrafine grains at least.
This technique is practically suitable for achieving ultrafine graining of about 1 to 3 μm, and has a fracture surface transition temperature (vTrS) of 2 mm V notch Charpy impact test of about −100 ° C. to −120 ° C. Although it is a sufficient technique for obtaining toughness stably, in order to achieve further low temperature toughness, it is necessary to further refine the crystal grain size.
[0004]
At present, ultrafine grain refinement on the order of 1 μm has been achieved by ultrafine graining technology similar to the strong machining to α and the subsequent machining and recrystallization of alpha as already disclosed by the present inventors. It's getting on. For example, the Japan Iron and Steel Institute 135th Spring Lecture Meeting Proceedings (CAMP-ISIJ, vol.11, No.3, 1988), page 566, by a method by other pass warm working with the deformation direction changed. An α particle size of about 1 μm has been achieved. However, both are limited to laboratory studies and have not yet been established as practical and industrial means.
[0005]
[Problems to be solved by the invention]
The present invention further develops the ultrafine graining technology by processing and recrystallization that the present inventors have already disclosed, without requiring special equipment, complicated hot working or heat treatment process, And, without limiting to the chemical composition in a special narrow range, it achieves an ultrafine α structure with an average α particle size of 1 μm or less over the entire thickness of the steel, and has extremely excellent toughness with a vTrS of −150 ° C. or less. The present invention provides a method for producing a high-tensile steel material having a tensile strength of about 500 MPa to 950 MPa.
[0006]
[Means for Solving the Problems]
According to the results of the present inventors' research so far, the dominant influence on the crystal grain size obtained by processing and recrystallization of α in the hot rolling stage is from the γ / α two-phase region to α. It is the cumulative reduction ratio in the region and the fineness of the structure before processing. Conventionally, as long as α is ultrafine-grained, it has been considered that α must be sufficiently generated by transformation before being processed into α. In other words, ultra-fine particles can be achieved in steels with relatively low hardenability, or by controlling the cooling so that agglomerate α is finely produced and processing with a large cumulative rolling reduction is applied to the α main structure. did. However, depending on the method, the α particle size usually obtained is in the range of 1 μm to 3 μm, and a crystal particle size of 1 μm or less can be obtained according to a very extreme heat history or severe processing conditions. It has also been found that it is not easy to stably achieve 1 μm or less.
[0007]
Therefore, as a result of further detailed research, the inventors of the present invention naturally have a large influence on the refinement of the previous structure and hot working on the crystal grain size. In order to stably obtain an ultrafine grain structure of 1 μm or less, a fine bainite transformed from unrecrystallized γ containing a large amount of dislocations controlled and rolled at a lower temperature than a fine massive α-based structure. It has been found that it is more advantageous to recrystallize the structure mainly composed of
[0008]
Furthermore, in addition to the above findings, it is necessary to increase the hardenability of the steel more than a certain amount in order to surely cause the bainite transformation, not the bulk α even from fine unrecrystallized γ, a structure different from α In order to recrystallize and make ultrafine grained bainite, it is necessary to set processing conditions different from those in the case where the previous structure is bulk α, and it is also found that an ultrafine grain structure of 1 μm or less is obtained. It led to a new invention. The summary is as follows.
[0009]
(1)% By mass
C: 0.01 to 0.25%
Si: 0.03-1.0%
Mn: 0.30 to 3.0%
Al: 0.003-0.1%
N: 0.001 to 0.01%
As impurities
P: 0.01% or less
S: 0.005% or less,
Contains
Cr: 0.01 to 2.0%
Ni: 0.01-6.0%
Mo: 0.01 to 2.0%
Cu: 0.01 to 1.5%
Of these, it contains one or more, consists of the remainder Fe and inevitable impurities, and(Formula 1) FormulaThe ideal quench critical diameter (DI)But,25.4-762mmIn the hot rolling of the steel slab,
Ac3After heating to a temperature above the transformation point and below 1150 ° C,
After performing the first rolling with a cumulative rolling reduction of 55 to 80%, a rolling start temperature of 850 ° C. or lower, and an end temperature of 750 ° C. or higher,
Second rolling is performed in which rolling with a rolling reduction rate of 30 to 70% per pass includes one or more passes, the cumulative rolling reduction rate is 50 to 95%, the rolling start temperature is 700 ° C. or lower, and the end temperature is 600 ° C. or higher. It is characterized byLowA method for producing high-tensile steel with excellent thermal toughness.
D I (Mm) = 12.7・ (C%) 0.5 ・ (1 + 0.64 ・ Si%)
・ (1 + 4.10 ・ Mn%) ・ (1 + 0.27 ・ Cu%)
・ (1 + 0.52 ・ Ni%) ・ (1 + 2.33 ・ Cr%)
・ (1 + 3.14 ・ Mo%) ・ (1 + 1.60 ・ W%)
...(Formula 1)
[0010]
(2) The aboveFirst rollingAs a pretreatment of steel,3Heating to a temperature not lower than the transformation point and not higher than 1050 ° C., and cooling to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s,As described in (1)A method for producing high-tensile steel with excellent low-temperature toughness.
[0011]
(3) The aboveFirst rollingAs a pretreatment of steel,3After heating to a temperature not lower than the transformation point and not higher than 1150 ° C., preliminary hot rolling with a cumulative reduction rate of 20 to 50% is performed, and cooling to 300 ° C. or lower is performed at a cooling rate of 0.2 to 20 ° C./s. The low temperature toughness described in (1) aboveExcellent high strength steelManufacturing method.
[0012]
(4)% By mass
C: 0.01 to 0.25%
Si: 0.03-1.0%
Mn: 0.30 to 3.0%
Al: 0.003-0.1%
N: 0.001 to 0.01%
As impurities
P: 0.01% or less
S: 0.005% or less,
Contains
Cr: 0.01 to 2.0%
Ni: 0.01-6.0%
Mo: 0.01 to 2.0%
Cu: 0.01 to 1.5%
Among them, the ideal quenching critical diameter (D) containing one or more of them, consisting of the balance Fe and inevitable impurities, and represented by the following formula (formula 1) I ) Is prior to hot rolling of steel slabs of 25.4 to 762 mm,
The steel slab is made of Ac 3 Heat to a temperature not lower than the transformation point and not higher than 1050 ° C. and cool to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s.After pretreatment,
After heating in a two-phase region with an austenite fraction of 50-90%, the rolling reduction rate per pass includes 30-70% or more rolling, the cumulative reduction rate is 50-95%, the rolling start temperature is A method for producing a high-tensile steel material excellent in low-temperature toughness, comprising rolling at 700 ° C. or lower and an end temperature of 600 ° C. or higher.
D I (Mm) = 12.7 · (C%) 0.5 · (1 + 0.64 · Si%)
・ (1 + 4.10 ・ Mn%) ・ (1 + 0.27 ・ Cu%)
・ (1 + 0.52 ・ Ni%) ・ (1 + 2.33 ・ Cr%)
・ (1 + 3.14 ・ Mo%) ・ (1 + 1.60 ・ W%)
...(Formula 1)
(5) Instead of the pretreatment, the steel piece is replaced with Ac 3 Pre-treatment of heating to a temperature not lower than the transformation point and not higher than 1150 ° C., followed by preliminary hot rolling with a cumulative reduction rate of 20 to 50% and cooling to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s. The method for producing a high-strength steel material having excellent low-temperature toughness as described in (4) above.
[0013]
(6)The steel material after the end of the last hot rolling is accelerated to 20 to 500 ° C. at a cooling rate of 2 to 40 ° C./s.Before(1) ~(5)Low temperature toughnessExcellent high strength steelManufacturing method.
[0014]
(7)450 ° C or higher, Ac1Tempering below the transformation point, Described in (6)Low temperature toughnessExcellent high strength steelManufacturing method.
[0015]
(8) By mass%
Ti: 0.003-0.10%
V: 0.005-0.50%
Nb: 0.003-0.10%
Zr: 0.003-0.10%
Ta: 0.005 to 0.20%
W: 0.01 to 2.0%
B: 0.0003 to 0.0020%
It is characterized by further containing one or more ofThe above(1) ~(7)Low temperature toughnessExcellent high strength steelManufacturing method.
[0016]
(9) In mass%,
Mg: 0.0005 to 0.01%
Ca: 0.0005 to 0.01%
REM: 0.005-0.10%
It is characterized by further containing one or more of themThe above(1) ~(8)Low temperature toughnessExcellent high strength steelManufacturing method.
[0017]
DETAILED DESCRIPTION OF THE INVENTION
The present invention is described in detail below.
The present invention is characterized in that processing is applied to a bainite main structure as a new means for surpassing the conventional achievement level in a method of ultrafine graining α by processing and recrystallization. Since the bainite structure has a lath structure, it is much more advantageous than the bulk α structure in terms of ease of recrystallization and the fineness of the structure in terms of controlling the grain size after recrystallization. However, in order to make the final α grain size 1 μm or less, it is not sufficient to simply use a bainite structure, but a bainite transformed from unrecrystallized γ having a fine grain size and high dislocation density. We found that this was necessary through detailed experiments.
[0018]
In order to stably transform bainite from fine-grained non-recrystallized γ, it is necessary to optimize the composition of the steel as well as to limit the individual elements for securing strength and toughness and to be an index of hardenability. Hardening critical diameter (DIIt is also important to keep a certain value above. That is, an optimal combination of chemical composition and manufacturing conditions is required. Hereinafter, the detailed description of the present invention will be divided into manufacturing conditions and chemical composition.
[0019]
First, the manufacturing conditions of the present invention will be described.
The method of the present invention firstly states that “the individual elements have the proper chemical composition, and(Formula 1) FormulaA steel slab with an appropriate DI value3After heating to a temperature not lower than the transformation point and not higher than 1150 ° C., after performing the first rolling at a cumulative rolling reduction of 55 to 80%, a rolling start temperature of 850 ° C. or lower and an end temperature of 750 ° C. or higher, one pass Perform a second rolling with a rolling reduction rate of 30-70% per pass, one or more passes, a cumulative reduction rate of 50-95%, a rolling start temperature of 700 ° C. or lower, and an end temperature of 600 ° C. or higher. The first wayAndFurthermore, in order to achieve further ultra-fine graining and ensure uniformity,3Heat to a temperature not lower than the transformation point and not higher than 1050 ° C., and cool to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s.3After heating to a temperature not lower than the transformation point and not higher than 1150 ° C., rolling is performed at a cumulative reduction rate of 20 to 50% and cooling to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s. The method of performing the first method after performing the processing is the second method.AndFurthermore, on the premise that the pre-structure refinement treatment is performed, “the steel slab is reheated to a two-phase region having an austenite fraction of 50 to 90%, and the rolling reduction per pass is 30 to 70%. A method that includes rolling at least one pass, performs rolling with a cumulative rolling reduction of 50 to 95%, a rolling start temperature of 700 ° C. or lower, and an end temperature of 600 ° C. or higher is referred to as a third method.
[0020]
Note that the steel slab in the present invention is an ingot obtained by solidifying molten steel with a mold, a slab such as a slab, billet, or bloom produced by a continuous casting method, and these slabs, mainly for the purpose of adjusting the shape and reducing segregation. It refers to all slabs, steel slabs, and steel plates before being subjected to hot rolling and heat treatment, which are requirements of the present invention, such as steel slabs obtained by partial rolling and hot rolling.
[0021]
The three methods described above are the basic requirements of the present invention. However, according to the necessity of material adjustment such as strength and toughness, after completion of the hot rolling, “20 to 20 ° C./s at a cooling rate of 2 to 40 ° C./s”. Accelerated cooling from ℃ to 300 ℃ ”or“ 450 ℃ or higher, Ac1It is also possible to perform tempering below the transformation point.
[0022]
The three methods that are the basic requirements of the present invention will be described in more detail.
First, the first method will be described. The steel slab is heated prior to hot rolling.ThreeIt is necessary to set the temperature within the range from the transformation point to 1150 ° C. This is because the heating temperature is AcThreeIf it is less than the transformation point, the part that is not gamma becomes a coarse alpha structure, and is not ultrafine-grained even by subsequent hot rolling, so that the final structure becomes noticeably mixed grains, conversely When the temperature is higher than 1150 ° C., the heated γ particle size becomes coarse, the γ before the transformation becomes insufficiently refined, and the final ultrafine grain becomes insufficient.
[0023]
Steel pieces are AcThreeAfter heating at the transformation point or higher and 1150 ° C. or lower, the first rolling with a cumulative rolling reduction of 55 to 80%, a rolling start temperature of 850 ° C. or lower and an end temperature of 750 ° C. or higher is performed, and γ grains before transformation Adjust the diameter and form. That is, the cumulative rolling reduction needs to be 55% or more in order to make γ finer and sufficiently introduce processing strain. As the cumulative rolling reduction is larger, it is advantageous for the introduction of finer grains and dislocations. However, if the cumulative rolling reduction in the rolling is excessive, it is difficult to secure the rolling reduction in the subsequent second rolling, The upper limit in the first rolling is 80% at which the rolling effect is saturated. On the other hand, if the cumulative rolling reduction is less than 55%, the accumulation of strain in unrecrystallized γ is insufficient, the subdivision of the bainite main structure after transformation becomes insufficient, and the final crystal grain size is refined. Will be insufficient. The first rolling temperature is limited to a start temperature of 850 ° C. or less and an end temperature of 750 ° C. This is because rolling above 850 ° C. has little effect on the γ grain refinement and the introduction of dislocations. This is because it is advantageous to set the rolling start temperature to 850 ° C. or lower in order to fully exhibit the effect of No. 2, and when the end temperature is less than 750 ° C., the rolling effect is saturated while the second rolling is performed. This is because it is difficult to ensure the temperature.
[0024]
If the rolling temperature is 850 ° C. to 750 ° C. within the range of the present invention, the rolling effect of each pass is almost accumulated, so there is no need to limit the rolling reduction of each pass, and the effect of the first rolling is the cumulative rolling reduction. Is achieved if it satisfies the scope of the present invention.
In addition, in this invention, it does not prevent applying the rolling mainly for shape adjustment etc. at the temperature of more than 850 degreeC before this rolling. That is, as long as the cumulative reduction ratio in the first and second rolling can be ensured and the effect of the present invention is not impaired, processing at a temperature higher than 850 ° C. may be added.
[0025]
After adjusting the grain size and form of γ before transformation by the first rolling to refine the bainite main structure after transformation, rolling with a rolling reduction rate of 30 to 70% per pass is 1 pass. Including the above, the second rolling is performed in which the cumulative rolling reduction is 50 to 95%, the rolling start temperature is 700 ° C. or lower, and the end temperature is 600 ° C. or higher.
By the second rolling, the bainite main structure is processed, and through the introduction / recovery / recrystallization process of dislocations, not only α but also the bainite structure is subdivided into an ultrafine grain structure of α as a whole. For that purpose, first, it is necessary to make bainite structure at least 30% or more before sufficiently transforming γ and entering the second rolling. The remainder may be α, γ, or other tissue. If the first rolling is performed, not only bainite but also other transformation structures are sufficiently refined, so there is no problem in ultrafine graining. Further, even if untransformed γ is present, it is transformed into bainite during rolling, and this is not a problem. However, in order to ensure the microstructure, the rolling start temperature needs to be 700 ° C. or lower. If it exceeds 700 ° C., the above pre-processed morphology cannot be stably guaranteed over the entire chemical composition of the present invention. Moreover, the rolling end temperature needs to be 600 degreeC or more. This is because if it is less than 600 ° C., recrystallization of the bainite structure becomes difficult and there are many cases where ultrafine graining is not sufficient, and deformation resistance becomes excessive and the load on the rolling mill becomes excessive. This is because it is difficult to ensure the shape of the steel material.
[0026]
For the above reason, the temperature of the second rolling needs to be in the range of 700 ° C. to 600 ° C., but it is also necessary to define the cumulative rolling reduction and the rolling amount of each pass. That is, in order to sufficiently recrystallize the bainite main structure in rolling at 700 ° C. to 600 ° C. to make it ultrafine grained, it is first assumed that the cumulative rolling reduction is 50 to 95%. That is, if the cumulative rolling reduction is less than 50%, recrystallization is difficult regardless of the rolling conditions of each pass, and there is a limit to the crystal grain size that can be achieved even if recrystallization is performed, so the cumulative rolling reduction is 50 % Or more is required. The larger the cumulative rolling reduction, the more advantageous for finer graining. However, if the cumulative rolling reduction ratio exceeds 95%, the finer graining effect is saturated, while problems arise in securing the temperature and productivity. Therefore, in the present invention, the upper limit is 95%. It was.
[0027]
In the second rolling, the rolling temperature and the cumulative rolling reduction are limited for the above reasons. However, in order to process the bainite main structure so that the α grain size is 1 μm or less, only the rolling temperature and the cumulative rolling reduction are limited. Is not sufficient, and when performing one-pass or multi-pass rolling with a cumulative rolling reduction of 50 to 95%, it is necessary to include one or more rollings with a rolling reduction of 30 to 70% per pass. In the present invention, the rolling temperature is considerably lower due to the need to process the bainite main structure in this invention, but in that case, recrystallization is more difficult than when processing is performed on the α structure main body. It is necessary to promote recrystallization by applying processing under large pressure in such a pass. Also, the difficulty of recrystallization varies depending on the crystal orientation, and depending on the orientation, simply increasing the cumulative rolling reduction does not recrystallize to the end, so it tends to be a mixed grain structure, but also to eliminate such a mixed grain structure It is preferable to perform rolling under large rolling at any stage of rolling. Therefore, the lower limit rolling reduction is 30%. In this case, when the processing rate is large and the temperature rises due to processing heat generation, the recovery and recrystallization of the structure is promoted, and the structure is uniformly refined.
[0028]
The larger the rolling reduction, the more preferable for the fine graining,
(1)Recrystallized α grains grow and become coarse due to excessive processing heat generation.
(2)If the rolling reduction of one pass is excessive, the rolling mill is overloaded.
(3)It adversely affects the steel plate shape.
For these three reasons, the upper limit of the rolling reduction per pass is 70%. Rolling with a 1-pass rolling reduction of 30 to 70% is more effective at the initial stage of rolling for the purpose of making the α grain size finer, and the cumulative rolling reduction and rolling temperature range are within the scope of the present invention. Then, the necessary ultrafine grain structure can be obtained at any stage of rolling. Moreover, although it is preferable that there are many passes under 30-70% reduction, if at least one pass is included, there is no problem in ultrafine-graining.
[0029]
That is, “Each element has an appropriate chemical composition, and(Formula 1) FormulaDIAc3After heating to a temperature not lower than the transformation point and not higher than 1150 ° C., after performing the first rolling at a cumulative rolling reduction of 55 to 80%, a rolling start temperature of 850 ° C. or lower and an end temperature of 750 ° C. or higher, one pass Perform a second rolling with a rolling reduction rate of 30-70% per pass, one or more passes, a cumulative reduction rate of 50-95%, a rolling start temperature of 700 ° C. or lower, and an end temperature of 600 ° C. or higher. This is the reason for limiting the first method characterized by the above. Next, the reason for limitation of the third method is shown together with the reason for the pre-structural refinement process. The pre-structure refinement process is also effective in the first method, and does not hinder the effect of the first method. By performing the pre-structure refinement process before rolling by the first method, Effective for obtaining toughness higher than that of the first method by reducing the size of the fine particles and the mixed particle size.WaySecond method.
[0030]
The third method is “the preprocessing according to claim 2 or 3.Same preprocessing asAfter refining the previous structure byWithout performing the first rolling as in the first method,Reheating to a two-phase region with an austenite fraction of 50-80%, including one or more passes with a rolling reduction of 30-70% per pass, a cumulative reduction of 50-95%, and a rolling start temperature of Rolling is performed at 700 ° C. or lower and an end temperature of 600 ° C. or higher ”. That is, unlike the first method, heating is performed to a two-phase region intermediate between the γ region and the α region before hot rolling.The same pretreatment as described inWhen refined, as in the present invention, as in the present invention, in a hardenable steel that has a bainite main structure with fine and non-recrystallized γ, reheating directly into the two-phase region is more effective during rolling. This is because the bainite transformation is promoted from the recrystallized γ, and a desired structure can be stably formed, and the temperature waiting time can be shortened by the lower heating temperature, leading to an improvement in productivity.
[0031]
When reheating to a two-phase region after refining the previous structure, it is necessary to control the γ fraction within a certain range because of the necessity of using a bainite-based structure in the present invention. That is, when the γ fraction is small, it becomes an α-based structure, and it is difficult to obtain an ultrafine grain structure of 1 μm or less. The lower limit of the γ fraction for making the bainite main structure necessary for setting the α particle size of the final structure to 1 μm or less is determined to be 50% based on experiments. On the other hand, if the γ fraction is too large, it is substantially the same as the γ single phase. Therefore, the γ fraction having the effect of promoting bainite transformation has an upper limit of 90%.
[0032]
As described above, the preprocessing according to claim 2 or 3Same preprocessing asThe rolling condition after the previous structure is refined by the first method is the first methodSecond rolling ofIs exactly the same.
Next, in the first method, further ultrafine grainingAlthough it is a process that is adopted as necessary for the purpose,In the third method, a pre-process for pre-structural refinement, which is an essential process for ultrafine graining, will be described.
[0033]
Claim 2Or 4Pretreatment described inThe steelThe coarse solidified structure of the piece is eliminated, and the structure before the processing in the two-phase region to the α region for ultrafine graining is refined. At that time, the steel piece3Heating to a temperature not lower than the transformation point and not higher than 1050 ° C.3If it is less than the transformation point, a coarse structure remains regardless of the subsequent cooling conditions, and if it exceeds 1050 ° C., the γ particle size during heating becomes coarse and the structure after cooling becomes insufficiently refined. It is. Moreover, it is necessary to control the cooling after heating to 300 ° C. or less at a cooling rate of 0.2 to 20 ° C./s. This condition is necessary to refine the transformation structure by subsequent cooling even if the heating temperature is specified as described above and the heating γ grain size is refined.
[0034]
That is, since the steel slab is generally thick, the cooling inside the steel slab is often gradually cooled even if it is allowed to cool, in which case the transformation structure becomes coarse. Therefore, in order to refine the structure after cooling, it is necessary to cool at 0.2 ° C./s or more. A larger cooling rate is preferable for refining the structure, but it is not practical to make the cooling rate extremely large for steel slabs with a large plate thickness. Since the appearance of an extremely coarse structure harmful to the crystal structure can be suppressed, the upper limit is set to 20 ° C./s in the present invention. Moreover, the appearance of such a coarse structure can be suppressed regardless of the subsequent cooling conditions if the controlled cooling is performed up to 300 ° C. If the steel piece thickness is relatively small and cooling is performed at 0.2 ° C./s or more by air cooling, naturally air cooling may be used.
[0035]
Claim 2 aboveOr in the pre-processing described in 4The structure is refined only by heat treatment.Or in the pretreatment described in 5The steel slab structure is refined by hot rolling. Although not only heat treatment but also processing by rolling is advantageous for refining the structure, it also reduces the thickness of the steel slab, so the thickness of the steel slab is thicker than the finished sheet thickness and the cumulative rolling reduction is small. It can be said that this method is preferable when the degree is high.
[0036]
In the method, first, the steel slab is made Ac.3Reheat to a temperature above the transformation point and below 1150 ° C. Ac3Reheating above the transformation point is claimed in claim 2.Or in case of pre-processing described in 4On the other hand, since the subsequent rolling has the effect of refining γ, the upper limit of the reheating temperature is relaxed compared to the case without rolling, but the reheated γ particle size is excessive. In the present invention, the upper limit is set to 1150 ° C., because the grain refinement becomes insufficient even by subsequent rolling.
[0037]
Ac3After reheating to a temperature not lower than the transformation point and not higher than 1150 ° C., hot rolling is performed to refine the γ structure. At this time, the cumulative rolling reduction of the hot rolling is set to 20 to 50%. This is Ac3In the case of rolling following reheating at a transformation point to 1150 ° C., if the cumulative rolling reduction is less than 20%, the γ grains are not sufficiently refined and the resulting transformation structure is not sufficiently refined even when the corner rolling is performed. On the other hand, if it exceeds 50%, the cumulative rolling reduction at the rolling stage for the final ultrafine graining cannot be sufficiently obtained. Claim 2 for cooling after hot rollingOr the pretreatment described in 4In the same manner as described above, as a condition for sufficiently miniaturizing the transformation structure, cooling is performed to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s.
[0038]
The above is an explanation of the basic requirements in the present invention, but other reasons for limiting the requirements related to the manufacturing method are added below.
That is, for the purpose of adjusting the strength of the steel material and improving the toughness, the steel material after the final rolling can be accelerated and cooled to 20 ° C. to 500 ° C. at a cooling rate of 2 to 40 ° C./s. This is because the internal structure is refined by accelerated cooling and the second phase is changed to a harder one. In order to exhibit this effect, the cooling rate needs to be 2 ° C./s or more. A higher cooling rate is advantageous for microstructure refinement and hard phase formation, but it is practically difficult to increase the cooling rate indefinitely with thick materials, so that the accelerated cooling effect on the structure control is not saturated. In the present invention, the upper limit is 40 ° C./s. Moreover, it is preferable to perform this accelerated cooling to 20 to 500 degreeC. In other words, cooling to below 20 ° C is not effective for controlling the structure, and if the accelerated cooling stop temperature exceeds 500 ° C, the structure is not sufficiently refined and it is difficult to form a hard phase. This is because the toughness may deteriorate. In addition, according to the accelerated cooling condition defined in the present invention, it is also effective for suppressing the grain growth of the ultrafine grain structure of the surface layer once formed.
[0039]
Further, tempering treatment can be performed for the purpose of strength adjustment, toughness improvement, and shape improvement. In that case, it is an essential requirement not to impair the ultrafine grain structure formed in the surface layer portion. In the present invention, the tempering temperature is set to 450 ° C. to Ac.1Although it is limited to the transformation point, the effect of tempering is not clear at less than 450 ° C.1This is because if the transformation point is exceeded, the morphology of the superfine grain structure in the surface layer portion is impaired. However, it is more preferable that the tempering temperature does not exceed 700 ° C. in order to more reliably suppress the grain growth of the ultrafine-grained layer. In addition, within the tempering temperature range of the present invention, the heating and holding time for tempering is arbitrary, but similarly, from the viewpoint of preserving the ultrafine grain structure of the surface layer portion, the holding time isWithin 5 hoursIt is preferable that
[0040]
The above are the requirements for the method for producing a high-strength steel material excellent in low-temperature toughness according to the present invention. In order to exert the effect by the production method, it is necessary to limit each chemical component for the reasons described below. There is.
[0041]
That is, C is contained as an effective component for improving the strength of steel, and if it is less than 0.01%, it is difficult to ensure the strength necessary for structural steel, but an excess content exceeding 0.25%. Lowers the toughness and weld crack resistance of the base metal and the welded portion, so the content was made 0.01 to 0.25%.
[0042]
Next, Si is an element effective as a deoxidizing element and for securing the strength of the base material. However, if it is less than 0.03%, deoxidation is insufficient and it is disadvantageous for securing the strength. On the contrary, an excessive content exceeding 1.0% forms a coarse oxide and causes deterioration of ductility and toughness. Therefore, the range of Si is set to 0.03 to 1.0%.
[0043]
Further, Mn is an element necessary for ensuring the strength and toughness of the base material, and it is necessary to contain at least 0.3% or more, but if it is contained excessively, it may cause hard phase formation or grain boundary embrittlement. In order to deteriorate the toughness of the base metal, the toughness of the welded portion, and the weld cracking property, the upper limit was made 3.0% within the allowable range of the material.
[0044]
Al is an element effective for deoxidation, austenite grain size reduction, etc., but in order to exert the effect, it is necessary to contain 0.003% or more. On the other hand, if it exceeds 0.1% and excessively contained, a coarse oxide is formed and the ductility is extremely deteriorated.
[0045]
N is combined with Al and Ti and effectively works for refining γ grains. Further, it is impossible to remove N in steel completely industrially, and reducing it more than necessary is not preferable because it places an excessive load on the manufacturing process. Therefore, the lower limit is set to 0.001% as a range that can be industrially controlled and the load on the manufacturing process is allowable. If excessively contained, solid solution N increases, which may adversely affect ductility and toughness, so the upper limit is made 0.01% as an acceptable range.
[0046]
In the present invention, in order to reliably transform a fine bainite structure from fine / unrecrystallized γ, a critical quench critical diameter (DI) Must be made appropriate, but the desired D can be obtained only by C, Si, Mn without affecting other characteristics.ISince it is difficult to adjust to a value, it is necessary to form a bainite structure by adding one or more of Cr, Ni, Mo, and Cu, which greatly contribute to hardenability. However, it is necessary to limit each addition amount as follows.
[0047]
Both Cr and Mo are effective elements for improving the strength of the base material, but in order to produce a clear effect, 0.01% or more is necessary. On the other hand, if added over 2.0%, toughness In addition, since the weldability tends to deteriorate, the range is 0.01 to 2.0%.
[0048]
Ni is a very effective element that can improve the strength and toughness of the base material at the same time. However, in order to exert the effect, Ni needs to be contained in an amount of 0.01% or more. If the content increases, the strength and toughness are improved, but even if added over 6.0%, the effect is saturated, but the weldability deteriorates, so the upper limit is made 6.0%. In addition, if the addition exceeds 6.0%, DIEven if the value is within the range of the present invention, the bainite transformation is suppressed, and an undesirable effect for ultrafine graining is also manifested.
[0049]
Next, Cu has substantially the same effect as Ni, but if it exceeds 1.5%, there is a problem in hot workability, so it is limited to a range of 0.01 to 1.5%.
[0050]
P and S are impurity elements and elements that degrade ductility and toughness, and it is preferable to reduce them as much as possible. However, material deterioration is not so great that the upper limit of P is 0.01% and the upper limit of S is acceptable. Is limited to 0.005%.
[0051]
The above is the reason for limiting the basic components of the steel material of the present invention. In the present invention, one of Ti, V, Nb, Zr, Ta, W, and B is selected as necessary for the adjustment of strength and toughness. It can contain seeds or two or more.
[0052]
Ti is an element that contributes to improving the strength of the base metal by precipitation strengthening, and is also an element effective for refinement of the heated austenite grain size by the formation of TiN, and is also an element effective for improving toughness. Needs to be contained in an amount of 0.003% or more. On the other hand, if it exceeds 0.10%, coarse precipitates and inclusions are formed to deteriorate toughness and ductility, so the upper limit is made 0.10%.
[0053]
V is also an element effective for improving the strength by forming VN. However, if contained excessively, the toughness deteriorates due to precipitation embrittlement. Accordingly, the range in which the effect can be exhibited without causing significant deterioration in toughness is limited to the range of 0.005 to 0.50%.
[0054]
Nb is an element effective for improving strength and toughness by forming Nb (C, N), but if contained excessively, toughness deteriorates due to precipitation embrittlement. Therefore, the range in which the effect can be exhibited without causing deterioration of toughness is limited to the range of 0.003 to 0.10%.
[0055]
Zr is also an element forming a nitride and has the same effect as Ti, but 0.003% or more is required to exhibit the effect. On the other hand, if it exceeds 0.10%, as in Ti, coarse precipitates and inclusions are formed to deteriorate toughness and ductility, so the content is limited to the range of 0.003 to 0.10%.
[0056]
Ta is also an element effective in improving strength and toughness, but 0.005% or more is necessary to exert the effect. On the other hand, if it exceeds 0.20%, precipitation embrittlement, coarse precipitates, and toughness deterioration due to inclusions occur, so the upper limit is made 0.20%.
[0057]
W is effective in increasing the strength of the base metal by solid solution strengthening and precipitation strengthening, but 0.01% or more is necessary to exert the effect. On the other hand, if it exceeds 2.0% and contains excessively, toughness deterioration becomes remarkable, so the upper limit is made 2.0%.
[0058]
B is an element effective for improving toughness by solid solution N fixation and improving strength and toughness by improving hardenability because it binds to N in a small amount, but 0.0003% or more is necessary to exert the effect. is necessary. On the other hand, if the content exceeds 0.0020%, BN becomes coarse and adversely affects ductility and toughness. Moreover, in order to deteriorate weldability, the upper limit is made 0.0020%.
[0059]
Furthermore, in order to improve the base material toughness, the ductility, and the joint toughness, one or more of Mg, Ca, and REM can be contained as necessary.
Mg, Ca, and REM are all effective in improving ductility by suppressing the extension of sulfide during hot rolling. The oxide is refined to contribute to further ultrafine-graining, which is advantageous for improving the base material toughness, and also effectively improves the joint toughness. The lower limit content for exhibiting the effect is 0.0005% for Mg and Ca, and 0.005% for REM. On the other hand, if it is excessively contained, the sulfides and oxides are coarsened and the ductility and toughness are deteriorated. Therefore, the upper limits are 0.01% for Mg and Ca and 0.10% for REM, respectively.
[0060]
The above is the reason for the limitation of individual elements. In order to reliably transform the fine bainite structure from fine and non-recrystallized γ, not only the limitation of individual elements but also the optimization of hardenability The following is an index of hardenability(Formula 1) FormulaCritical quench critical diameter (DI)25.4-762mmIt is necessary to.
D I (Mm) = 12.7・ (C%) 0.5 ・ (1 + 0.64 ・ Si%)
・ (1 + 4.10 ・ Mn%) ・ (1 + 0.27 ・ Cu%)
・ (1 + 0.52 ・ Ni%) ・ (1 + 2.33 ・ Cr%)
・ (1 + 3.14 ・ Mo%) ・ (1 + 1.60 ・ W%)
...(Formula 1)
[0061]
(Formula 1) FormulaD indicated byIvalue25.4mmIf it is less than the range, the hardenability is insufficient, and the ratio of the bainite structure in the transformed structure becomes insufficient, making it difficult to obtain an ultrafine structure of 1 μm or less. On the other hand, DIvalue762mmIf the excess exceeds 1, the pre-structural refinement or two-phase region heating, etc., does not perform bainite transformation from γ, but martensite transformation, so the transformation temperature becomes as low as about 400 or less, transformation It becomes virtually impossible to process the tissue. For the above reason, in the present invention,(Formula 1) FormulaCritical quench critical diameter (DI)25.4-762mmLimited to.
[0062]
【Example】
The above is an explanation of the requirements of the present invention. Further, the effects of the present invention are shown based on examples.
Steel sheets were produced under the production conditions shown in Tables 2 and 3 using steel pieces having chemical compositions shown in Table 1. The characteristics of the ferrite structure (average diameter, mixed grain size, grain shape, etc. at the center of the steel sheet thickness) and mechanical properties (strength, 2 mm V notch Charpy impact characteristics, ESSO characteristics) of the manufactured steel sheet The measurement results are also shown in Tables 2 and 3. In addition, Table 2 is the result which showed this invention steel and the comparative example which relate to claims 1-3, Table 3 isClaims 4 and 5It is the result which showed this invention steel related to 1 and a comparative example.
[0063]
The observation of the structure and the measurement of the particle diameter are carried out based on the scanning electron microscope (SEM) structure photograph of the magnification 1000 times to 5000 times for the steels of the present invention and comparative examples with relatively fine ferrite particle diameters. The comparative example with a coarse particle size was carried out based on a photomicrograph of an optical microscope having a magnification of 500 times. Tensile properties were obtained by collecting round bar tensile test pieces from the center of the thickness in the direction perpendicular to the rolling direction (C direction). A 2 mm V notch Charpy impact test was also conducted by collecting a standard test piece having a thickness of 10 mm from the center of the thickness in the direction perpendicular to the rolling direction (C direction). The Charpy impact characteristics were evaluated by the fracture surface transition temperature (vTrS). Furthermore, the brittle crack propagation stopping property (arrestability) was investigated by a temperature gradient type ESSO test. For ESSO characteristics, specimens are taken from the C direction, and Kca is 400 kgf / mm.1.5 (= 3920 N / mm 1.5 )Was evaluated at a temperature (TKca 400).
[0064]
Test No. in Table 2 Test Nos. A1 to A9 and Table 3 A10 to A16 are steel plates produced by the production method of the present invention using steel numbers 1 to 11 having the chemical composition of the present invention, and all of them have an average ferrite particle size of 1 μm or less, and conventional ferrite processing -An ultrafine grain structure that is one level finer than the level of ultrafine grain steel by recrystallization is obtained, and extremely good toughness and arrestability are obtained. That is, the vTrS of the conventional ultrafine-grained steel having an average ferrite particle size of about 3 μm or less is about −100 to −150 ° C., whereas the vTrS of the ultrafine-grained steel of the present invention having an average particle size of 1 μm or less. Are all below -170 ° C, and most of them show a fracture mode mainly composed of ductile fracture even at liquid nitrogen temperature (-196 ° C). Reflecting such extremely high toughness, ESSO characteristics at full thickness However, it is apparent that TKca 400 is very good at −132 ° C. or lower, and the safety of the steel material is dramatically increased.
[0065]
On the other hand, from the results of Tables 2 and 3, the test No. deviating from the scope of the present invention. The steel plates B1 to B11 are the test Nos. Manufactured according to the present invention. It is clear that the toughness and arrestability are significantly inferior to the steel sheets of A1 to A16.
Test No. B1-B5 and No. B11 is an example in which sufficient toughness and arrestability could not be achieved although the production method satisfied the present invention because the chemical composition did not satisfy the present invention.
[0066]
That is, test no. B1(Formula 1) FormulaDISince the value is too small and the structure before the two-phase region processing does not become a bainite-based structure, the finally obtained ferrite grain size is 1.98 μm, which is not less than 1 μm, which is vTrS compared to the steel of the present invention. And arrestability (TKca 400) is inferior.
[0067]
Test No. Conversely, B2 is DISince the value is excessive and the hardenability is excessive, the first rolling is performed by the method according to claim 1 and the rolling is performed at a low temperature, but the majority is untransformed austenite. Therefore, an ultrafine-grained ferrite structure is not obtained, and therefore vTrS and arrestability (TKca 400) are significantly inferior to the steel of the present invention.
[0068]
In addition, Test No. B11 is Test No. Using the same steel piece as B2,Claim 5The two-phase region heating is applied, but the part that was an austenite phase during the two-phase region heating does not cause transformation before and during processing, and the final structure is elongated. As a result, the bainite-based organization was obtained. Like B2, ultrafine graining has not been achieved, and the toughness and arrestability are inferior to the steel of the present invention.
[0069]
Test No. In B3, since C is excessive, fine toughness is sufficient, but a good toughness level is not achieved.
Test No. B4 is inferior in both vTrS and TKca 400 compared to the steel of the present invention, although both P and S are excessive and ultrafine graining is sufficient.
Test No. In B5, since both Cr and Mo are excessive, a good toughness level is not achieved even though ultrafine graining is sufficient.
[0070]
On the other hand, test no. B6, B7, and B8 to B10 satisfy the present invention in chemical composition, but the manufacturing method deviates from the scope of the present invention. Is an example in which the characteristics are inferior to those of the steel sheet produced according to the present invention.
[0071]
Test No. B6 and B7 are comparative examples related to claim 1. In the second rolling, No. B6 has a rolling reduction of 30% or more in any rolling pass. In B7, since the cumulative rolling reduction ratio of the first rolling does not satisfy the range of the present invention, the final ferrite grain size is coarser than that of the steel sheet according to the present invention, and the toughness and arrestability are also the present invention. Inferior to steel.
[0072]
Tests B8-B10Claims 4 and 5It is a comparative example related to.
Test No. Since B8 is not subjected to the pre-structural refinement step before the two-phase region heating / rolling, a coarse structure remains in the two-phase region heating stage, and the final structure is also mixed / coarse, toughness, Arrest is inferior.
Test No. In B9, since rolling reduction after the two-phase region heating is not 30% or more in any rolling pass, ultrafine graining is not sufficient compared to the steel of the present invention, and both vTrS and TKca 400 It stays at the conventional level.
[0073]
Test No. B10 has a low austenite fraction in the heating stage due to the low temperature in the two-phase region, and as a result, is an essential requirement for achieving a ferrite grain size of 1 μm or less. Since the condition of “processing the microstructure mainly composed of fine bainite transformed from recrystallized γ” is not satisfied, the average ferrite grain size can only be reduced to about 2 μm, and by conventional processing and recrystallization of ferrite Only the toughness level of ultrafine-grained steel is obtained.
[0074]
From the above examples, according to the present invention, an ultrafine ferrite structure having an average ferrite grain size of 1 μm or less can be achieved over the entire thickness of the steel material, and production of a high-strength steel material having extremely good Charpy impact properties and arrestability It is clear that will be possible.
[0075]
[Table 1]
[Table 2]
[Table 3]
[Table 4]
[Table 5]
[Table 6]
[Table 7]
[Table 8]
【The invention's effect】
The present invention achieves an ultrafine ferrite structure with an average ferrite grain size of 1 μm or less over the entire thickness of the steel material, and has vTrS of −150 ° C. or less and extremely excellent low-temperature toughness and brittle crack propagation stopping properties. It provides a method for producing high-tensile steel materials with a tensile strength of about 500MPa to over 950MPa, and can improve the economics and productivity of high-tensile steel materials that can dramatically improve the safety of structures. It can be said that it is an invention with extremely high industrial value as a means that can be manufactured without loss.
Claims (9)
Ｃ ：０．０１〜０．２５％
Ｓｉ：０．０３〜１．０％
Ｍｎ：０．３０〜３．０％
Ａｌ：０．００３〜０．１％
Ｎ ：０．００１〜０．０１％
不純物として、
Ｐ ：０．０１％以下
Ｓ ：０．００５％以下、
を含有し、さらに
Ｃｒ：０．０１〜２．０％
Ｎｉ：０．０１〜６．０％
Ｍｏ：０．０１〜２．０％
Ｃｕ：０．０１〜１．５％
のうち、１種または２種以上を含有し、残部Ｆｅ及び不可避不純物からなり、かつ、下記（式１）式で示される理想焼入臨界直径（ＤＩ）が、２５．４〜７６２ｍｍである鋼片の熱間圧延にあたって、
Ａｃ３変態点以上、１１５０℃以下の温度に加熱した後、
累積圧下率が５５〜８０％で、圧延開始温度が８５０℃以下、終了温度が７５０℃以上の第１の圧延を行った後、
１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の第２の圧延を行うことを特徴とする、低温靭性に優れた高張力鋼材の製造方法。
Ｄ Ｉ （ｍｍ）＝１２．７・（Ｃ％）０．５・（１＋０．６４・Ｓｉ％）
・（１＋４．１０・Ｍｎ％）・（１＋０．２７・Ｃｕ％）
・（１＋０．５２・Ｎｉ％）・（１＋２．３３・Ｃｒ％）
・（１＋３．１４・Ｍｏ％）・（１＋１．６０・Ｗ％）
・・・ （式１） % By mass
C: 0.01 to 0.25%
Si: 0.03-1.0%
Mn: 0.30 to 3.0%
Al: 0.003-0.1%
N: 0.001 to 0.01%
As impurities
P: 0.01% or less S: 0.005% or less,
And further Cr: 0.01-2.0%
Ni: 0.01-6.0%
Mo: 0.01 to 2.0%
Cu: 0.01 to 1.5%
Among them, the ideal quench critical diameter (D I ) is 25.4 to 762 mm, containing one or two or more types, the balance being Fe and inevitable impurities, and represented by the following formula (Formula 1): In hot rolling of billets,
After heating to a temperature of Ac 3 transformation point or higher and 1150 ° C. or lower,
After performing the first rolling with a cumulative rolling reduction of 55 to 80%, a rolling start temperature of 850 ° C. or lower, and an end temperature of 750 ° C. or higher,
Second rolling is performed in which rolling with a rolling reduction rate of 30 to 70% per pass includes one or more passes, the cumulative rolling reduction rate is 50 to 95%, the rolling start temperature is 700 ° C. or lower, and the end temperature is 600 ° C. or higher. it shall be the said method for producing a high tensile steel having excellent low-temperature toughness.
D I (mm) = 12.7 · (C%) 0.5 · (1 + 0.64 · Si%)
・ (1 + 4.10 ・ Mn%) ・ (1 + 0.27 ・ Cu%)
・ (1 + 0.52 ・ Ni%) ・ (1 + 2.33 ・ Cr%)
・ (1 + 3.14 ・ Mo%) ・ (1 + 1.60 ・ W%)
... (Formula 1)
Ｃ ：０．０１〜０．２５％
Ｓｉ：０．０３〜１．０％
Ｍｎ：０．３０〜３．０％
Ａｌ：０．００３〜０．１％
Ｎ ：０．００１〜０．０１％
不純物として、
Ｐ ：０．０１％以下
Ｓ ：０．００５％以下、
を含有し、さらに
Ｃｒ：０．０１〜２．０％
Ｎｉ：０．０１〜６．０％
Ｍｏ：０．０１〜２．０％
Ｃｕ：０．０１〜１．５％
のうち、１種または２種以上を含有し、残部Ｆｅ及び不可避不純物からなり、かつ、下記（式１）式で示される理想焼入臨界直径（Ｄ Ｉ ）が、２５．４〜７６２ｍｍである鋼片の熱間圧延に先立ち、
該鋼片を、Ａｃ ３ 変態点以上、１０５０℃以下の温度に加熱し、０．２〜２０℃／ｓの冷却速度で３００℃以下まで冷却する前処理を施した後、
オーステナイト分率が５０〜９０％の二相域に加熱した後に、１パスあたりの圧下率が３０〜７０％の圧延を１パス以上含み、累積圧下率が５０〜９５％で、圧延開始温度が７００℃以下、終了温度が６００℃以上の圧延を行うことを特徴とする、低温靭性に優れた高張力鋼材の製造方法。
Ｄ Ｉ （ｍｍ）＝１２．７・（Ｃ％）０．５・（１＋０．６４・Ｓｉ％）
・（１＋４．１０・Ｍｎ％）・（１＋０．２７・Ｃｕ％）
・（１＋０．５２・Ｎｉ％）・（１＋２．３３・Ｃｒ％）
・（１＋３．１４・Ｍｏ％）・（１＋１．６０・Ｗ％）
・・・ （式１） % By mass
C: 0.01 to 0.25%
Si: 0.03-1.0%
Mn: 0.30 to 3.0%
Al: 0.003-0.1%
N: 0.001 to 0.01%
As impurities
P: 0.01% or less
S: 0.005% or less,
Contains
Cr: 0.01 to 2.0%
Ni: 0.01-6.0%
Mo: 0.01 to 2.0%
Cu: 0.01 to 1.5%
Among them, the ideal quench critical diameter (D I ) is 25.4 to 762 mm, which contains one or two or more, consists of the balance Fe and inevitable impurities, and represented by the following formula (Formula 1). Prior to hot rolling of billets,
The steel slab is heated to a temperature not lower than Ac 3 transformation point and not higher than 1050 ° C., and subjected to a pretreatment for cooling to 300 ° C. or lower at a cooling rate of 0.2 to 20 ° C./s ,
After heating in a two-phase region with an austenite fraction of 50-90%, the rolling reduction rate per pass includes 30-70% or more rolling, the cumulative reduction rate is 50-95%, the rolling start temperature is A method for producing a high-tensile steel material excellent in low-temperature toughness, comprising rolling at 700 ° C. or lower and an end temperature of 600 ° C. or higher.
D I (mm) = 12.7 · (C%) 0.5 · (1 + 0.64 · Si%)
・ (1 + 4.10 ・ Mn%) ・ (1 + 0.27 ・ Cu%)
・ (1 + 0.52 ・ Ni%) ・ (1 + 2.33 ・ Cr%)
・ (1 + 3.14 ・ Mo%) ・ (1 + 1.60 ・ W%)
... (Formula 1)
Ｔｉ：０．００３〜０．１０％
Ｖ ：０．００５〜０．５０％
Ｎｂ：０．００３〜０．１０％
Ｚｒ：０．００３〜０．１０％
Ｔａ：０．００５〜０．２０％
Ｗ ：０．０１〜２．０％
Ｂ ：０．０００３〜０．００２０％
の１種または２種以上をさらに含有することを特徴とする、請求項１〜７のいずれか１項に記載の低温靭性に優れた高張力鋼材の製造方法。 % By mass
Ti: 0.003-0.10%
V: 0.005-0.50%
Nb: 0.003-0.10%
Zr: 0.003-0.10%
Ta: 0.005 to 0.20%
W: 0.01 to 2.0%
B: 0.0003 to 0.0020%
1 or 2 types or more of these are further contained, The manufacturing method of the high strength steel materials excellent in the low temperature toughness of any one of Claims 1-7 characterized by the above-mentioned.
Ｍｇ：０．０００５〜０．０１％
Ｃａ：０．０００５〜０．０１％
ＲＥＭ：０．００５〜０．１０％
のうち１種または２種以上をさらに含有することを特徴とする、請求項１〜８のいずれか１項に記載の低温靭性に優れた高張力鋼材の製造方法。 % By mass
Mg: 0.0005 to 0.01%
Ca: 0.0005 to 0.01%
REM: 0.005-0.10%
The manufacturing method of the high-tensile steel material excellent in the low temperature toughness of any one of Claims 1-8 characterized by further containing 1 type, or 2 or more types.
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| CN107217201A (en) * | 2017-06-27 | 2017-09-29 | 包头钢铁（集团）有限责任公司 | A kind of marine drilling platform containing rare earth spud leg 600MPa seamless steel pipes and its production method |
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