JP3814112B2 - Super high strength steel pipe excellent in low temperature toughness of seam welded portion and manufacturing method thereof - Google Patents
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【０００１】
【発明の属する技術分野】
本発明は、天然ガス・原油輸送用ラインパイプとして広く使用でき、高圧化による輸送効率の向上及び外径・重量の低減による現地施工能率の向上が可能である９００ＭＰａ以上の引張強さを有するシーム溶接部の低温靱性の優れた超高強度ラインパイプおよびその製造方法に関するものである。
【０００２】
【従来の技術】
近年、原油・天然ガスの長距離輸送方法としてパイプラインの重要性がますます高まっている。現在、長距離輸送用の幹線ラインパイプとしては米国石油協会（ＡＰＩ）規格Ｘ６５が設計の基本になっており、実際の使用量も圧倒的に多い。しかし、(1) 高圧化による輸送効率の向上、(2) ラインパイプの外径・重量の低減による現地施工能率の向上、のためより高強度ラインパイプが要望されている。これまでにＸ８０（引張強さ６２０ＭＰａ以上）までのラインパイプの実用化がされているが、さらに高強度のラインパイプに対するニーズが強くなってきた。現在、超高強度ラインパイプ製造法の研究は、従来のＸ８０ラインパイプの製造技術（例えば、ＮＫＫ技報No.138(1992), pp24-31 およびThe 7th Offshore Mechanics and Arctic Engineering (1988), Volume V, pp179-185)を基本に検討されているが、これではせいぜい、Ｘ１００（引張強さ７６０ＭＰａ以上）ラインパイプの製造が限界と考えられる。Ｘ１００を越える超高強度ラインパイプについては、既に鋼板製造の研究は行われている（ＰＣＴ／ＪＰ９６／００１５５、００１５７）。しかし、このような超高強度ラインパイプでは従来のシーム溶接に関する技術は適用できず、シーム溶接技術に関する課題が解決できないと鋼板は製造できても鋼管の製造は不可能である。パイプラインの超高強度化は強度・低温靱性バランスを始めとして溶接熱影響部（ＨＡＺ）靱性、現地溶接性、継手軟化、バースト試験による管体破断など多くの問題を抱えており、これらを克服した画期的な超高強度ラインパイプ（Ｘ１００超）の早期開発が要望されている。
【０００３】
【発明が解決しようとする課題】
本発明は低温靱性のバランスが優れ、かつ現地溶接が容易な引張強さ（ＴＳ）９００ＭＰａ以上（ＡＰＩ規格Ｘ１００超）の超高強度ラインパイプ、特にバースト試験において溶接部破断がなく管体破断するシーム溶接部の低温靱性に優れた超高強度ラインパイプおよびその製造方法を提供するものである。
【０００４】
【課題を解決するための手段】
本発明の要旨は以下のとおりである。
（１）母材部の引張り強度が９００ＭＰａ以上であり、かつ、溶接金属部の引張り強度と母材部の引張り強度の差が−１００ＭＰａ以上であるシーム溶接鋼管であって、該シーム溶接鋼管の前記溶接金属部において、製管プロセスの鋼板付き合わせ部の仮付け溶接後に行われる本溶接によって形成される内面溶接金属部と外面溶接金属部の間隔が０ｍｍ超であり、かつ、内面溶接金属部と外面溶接金属部が前記仮付け溶接によって形成される仮付け溶接金属部とそれぞれ重複していることを特徴とするシーム溶接部の低温靱性の優れた超高強度溶接鋼管。
（２）重量％で、
Ｃ：０．０３〜０．１０％、
Ｓｉ：０．６％以下、
Ｍｎ：１．７〜２．５％、
Ｐ：０．０１５％以下、
Ｓ：０．００３％以下、
Ｎｉ：０．１〜１．０％、
Ｍｏ：０．１５〜０．６０％、
Ｎｂ：０．０１〜０．１０％、
Ｔｉ：０．００５〜０．０３０％、
Ａｌ：０．０６％以下、
を含有し、さらに重量％で、Ｂ：０．００３０％以下、Ｎ：０．００１〜０．００６％、Ｖ：０．１０％以下、Ｃｕ：１．０％以下、Ｃｒ：１．０％以下、Ｃａ：０．０１％以下、ＲＥＭ：０．０２％以下、Ｍｇ：０．００６％以下の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる母材部と、
重量％で、
Ｃ：０．０３〜０．１４％、
Ｓｉ：０．０５〜０４０％、
Ｍｎ：１．２〜２．２％、
Ｐ：０．０１０％以下、
Ｓ：０．０１０％以下、
Ｎｉ：１．３〜３．２％、
Ｃｒ、Ｍｏ、Ｖのうちの１種または２種以上の合計量が１．０〜２．５％、
Ｂ：０．００５％以下、
を含有し、残部が鉄および不可避的不純物からなる溶接金属部からなり、かつ、溶接金属部のＮｉ量が母材部のＮｉ量に比べて１％以上高く、溶接金属部部及び母材の溶接熱影響部を含むシーム溶接部の組織がベイナイト・マルテンサイトからなることを特徴とする上記（１）に記載のシーム溶接部の低温靱性の優れた超高強度溶接鋼管。
（３）重量％で、
Ｃ：０．０３〜０．１０％、
Ｓｉ：０．６％以下、
Ｍｎ：１．７〜２．５％、
Ｐ：０．０１５％以下、
Ｓ：０．００３％以下、
Ｎｉ：０．１〜１．０％、
Ｍｏ：０．１５〜０．６０％、
Ｎｂ：０．０１〜０．１０％、
Ｔｉ：０．００５〜０．０３０％、
Ａｌ：０．０６％以下、
を含有し、さらに重量％で、Ｂ：０．００３０％以下、Ｎ：０．００１〜０．００６％、Ｖ：０．１０％以下、Ｃｕ：１．０％以下、Ｃｒ：１．０％以下、Ｃａ：０．０１％以下、ＲＥＭ：０．０２％以下、Ｍｇ：０．００６％以下の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼板の両端部を付き合わせた後、該付き合わせ部を、重量％で、Ｃ：０．０１〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１．２〜２．４% を含有しＦｅを主成分とする溶接ワイヤーを用いて、外面から仮付け溶接を行った後、該仮付け溶接部を、重量％で、Ｃ：０．０１〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１．２〜２．４% 、Ｎｉ：４．０〜８．５％、Ｃｒ、Ｍｏ、Ｖの１種又は２種以上の合計量３．０〜５．０％を含有し、かつＮｉ量が前記鋼板のＮｉ量に比べて１％以上高いＦｅを主成分とする溶接ワイヤーおよびフラックスを用いて、溶接によって形成される内面溶接金属部と外面溶接金属部の間隔が０ｍｍ超であり、かつ、内面溶接金属部と外面溶接金属部が前記仮付け溶接によって形成される仮付け溶接金属部とそれぞれ重複するように、前記仮付け溶接部を内面及び外面から本溶接を行うことを特徴とするシーム溶接部の低温靱性の優れた超高強度溶接鋼管の製造方法。
（４）前記本溶接において、仮付け溶接部を内面及び外面からそれぞれ２パス以上の溶接を行うことを特徴とする上記（３）に記載のシーム溶接部の低温靱性の優れた超高強度溶接鋼管の製造方法。
（５）前記仮付け溶接として、ＭＡＧアーク溶接、ＭＩＧアーク溶接、ＴＩＧアーク溶接の何れか１つの方法を用い、前記本溶接として、サブマージアーク溶接、ＭＡＧアーク溶接、ＭＩＧアーク溶接、ＴＩＧアーク溶接の何れか１つの方法を用いることを特徴とする上記（３）または（４）のいずれかに記載のシーム溶接部の低温靱性の優れた超高強度溶接鋼管の製造方法。
【０００５】
【発明の実施の形態】
以下、本発明の内容について詳細に説明する。
本発明は９００ＭＰａ以上の引張強さ（ＴＳ）を有するシーム溶接部の低温靱性の優れた超高強度ラインパイプに関する発明である。この強度水準の超高強度ラインパイプでは、従来主流であるＸ６５と較べて約２倍の圧力に耐えるため、同じサイズで約２倍のガスを輸送することが可能になる。一方、Ｘ６５を用いて上記超高強度ラインパイプと同等なガス輸送効率を達成する場合は圧力を高めるために肉厚を厚くする必要があり、材料費、輸送費、現地溶接施工費が高くなってパイプライン敷設費が大幅に上昇する。これが９００ＭＰａ以上の引張強さ（ＴＳ）を有する低温靱性の優れた超高強度ラインパイプが必要とされる理由である。従来、このような引張強さが９００ＭＰａ以上の超高強度ラインパイプでは、極端に鋼管の製造が困難になるとともに鋼管の特に低温靱性の特性を確保することが困難になる。鋼管のシーム溶接部も含めた目標特性を保証するための目安として、バースト試験において溶接熱影響部及び溶接金属等で破断せずに管体での破断が達成されることとともにシーム溶接部の低温靱性を改善することが重要な技術的課題になる。従来の超高強度ラインパイプでは、溶接時にシーム溶接部の接熱影響部の会合部から１mmまでに旧オーステナイト粒界に沿って粗大なMA（Martensite-Austenite Constituent：マルテンサイトとオーステナイトの混成物）が生成しやすく、これが破壊の起点となり、吸収エネルギー値を著しく低下させる要因であった。したがって、従来の溶接熱影響部の１／２ｔ部の会合部あるいは会合部＋１mmにおけるV ノッチシャルピー吸収エネルギーは、−３０℃で５０Ｊ未満と低く、例えば−３０℃で６４Ｊ以上の目標を満足させることはかなり困難であった。
【０００６】
本発明者らは、引張強さが９００ＭＰａ以上の超高強度のラインパイプにおいてシーム溶接部の低温靱性を改善すべく、実験等により鋭意検討した。
図１及び図２に超高強度のラインパイプにおける従来の溶接部と本発明による溶接部を示す。通常鋼管の造管時のシーム溶接は、鋼板両端部を付き合わせた後、付き合わせ部を最初に外面からＭＡＧアーク溶接等で仮付け（以下仮付け溶接と言う）を行い、その後、その仮付け溶接部をさらに内面及び外面からサブマージドアーク溶接等で溶接（以下本溶接と言う）を行う。この際、従来の本溶接は、図１に示すように本溶接において外面からの溶接で形成された溶接金属部（以下外面溶接金属部という）と内面からの溶接で形成された溶接金属部（以下内面溶接金属部という）を互いに重複させるため、本溶接の前に行った仮付け溶接金属部が溶融・消失し、溶接入熱が過度に高くなり溶接熱影響部の旧オーステナイト粒が粗大化すると共に旧オーステナイト粒界に沿って生成するMAも粗大化し、これが溶接熱影響部のシャルピー吸収エネルギーを低下させ、また、溶接熱影響部の軟化がおこってバースト試験において溶接部からの破断（管体破断ではなく）を招く要因となることがわかった。
【０００７】
本発明者らは、従来のような外面溶接金属部と内面溶接金属部を互いに重複させた内外面からの本溶接における過度な溶接入熱の上昇に起因する溶接部低温靱性の低下という問題を解決するために、本溶接の最適条件について、詳細な検討をおこなった。その結果、図２に示すように、仮付け溶接後の本溶接において、外面溶接金属部と内面溶接金属部を重複させずに仮付け溶接金属部を溶融させず残存させ、本溶接時の過度な溶接入熱の上昇を避けることによって溶接熱影響部で粗粒部の旧オーステナイト粒径の粗大化及びMAの粗大化を抑制し、溶接熱影響部の１／２ｔ部の会合部あるいは会合＋１ｍｍでのＶノッチシャルピー吸収エネルギーを改善し、また、溶接熱影響部の軟化部が抑えられることによってバースト試験において管体破断（溶接部からの破断なし）が可能となることがわかった。また、このような本溶接を行う場合、溶接部の溶接欠陥の発生を防止するために、本溶接によって形成された内面溶接金属部及び外面溶接金属部のそれぞれと、その前の仮付け溶接によって形成された仮溶接金属部とを重複する必要があることがわかった。
【０００８】
以上の知見から、本発明では、引張強さが９００ＭＰａ以上の超高強度のラインパイプのシーム溶接部において、仮付け溶接後の本溶接によって形成される内面溶接金属部と外面溶接金属部の間隔（Δｄ）が０ｍｍ超（これに対して従来のΔｄは０ｍｍ以下の値となる）とし、かつ、内面溶接金属部及び外面溶接金属部が前記仮付け溶接によって形成された仮付け溶接金属部とそれぞれ重複することを要件とする。上記Δｄが０ｍｍ以下、つまり本溶接によって形成される内面溶接金属部と外面溶接金属部が重なると、上述のように本溶接の前に行った仮付け溶接金属部が溶融・消失し、溶接入熱が過度に高くなり、その結果、溶接熱影響部の旧オーステナイト粒が粗大化すると共に旧オーステナイト粒界に沿って生成するMAも粗大化し、溶接熱影響部のシャルピー吸収エネルギーの低下や溶接熱影響部の軟化が起こる。また、本溶接部によって形成される内面溶接金属部及び外面溶接金属部のそれぞれと、仮付け溶接によって形成される仮付け溶接金属部とが重複していなければ、溶接部の溶接欠陥が発生する。ここで本発明で最初に行う仮付け溶接の方法は、通常知られているＭＡＧアーク溶接でもＭＩＧアーク溶接でもＴＩＧ溶接でも良い。また、仮付け溶接の後に行う内外面溶接もサブマージアーク溶接でもＭＩＧアーク溶接でもＴＩＧ溶接でも良い。また、仮付け溶接後の本溶接における仮付け溶接部の内面及び外面からの溶接は、それぞれ１パス溶接であっても２パス以上の溶接であっても良いが、本溶接時の溶接入熱を出来る限り下げることで溶接熱影響部の軟化を抑えるという点から２パス以上の溶接がより好ましい。
【０００９】
また、本発明者らが上記溶接法で製造した鋼管を用いて多数のバースト試験を行った結果から、溶接金属の引張強度が、［母材部の強度］−１００（ＭＰａ）以上であれば溶接部から破断せず、管体から破断することがわかっている。従って、本発明では、溶接部の平均引張強度を［母材部の円周方向引張強度］−１００（ＭＰａ）以上とする。
【００１０】
次に本発明の鋼管を構成する母材部及び溶接金属部の成分および組織について説明する。
先ず、本発明の母材成分の限定理由は以下の通りである。
Ｃ量は、０．０３〜０．１０％に限定する。炭素は、鋼の強度向上に極めて有効であり、マルテンサイト組織において目標とする強度を得るためには、最低０．０３％は必要である。しかし、Ｃ量が多すぎると母材、ＨＡＺの低温靱性や現地溶接性の著しい劣化を招くので、その上限を０．１０％とした。さらに、望ましくは上限値は０．０８％が好ましい。
【００１１】
Ｓｉは脱酸や強度向上のために添加する元素であるが、多く添加するとＨＡＺ靱性、現地溶接性を著しく劣化させるので、上限を０．６％とした。鋼の脱酸はＡｌでもＴｉでも十分可能であり、Ｓｉは必ずしも添加する必要はない。
Ｍｎは本発明鋼のミクロ組織をマルテンサイト主体の組織とし、優れた強度・低温靱性のバランスを確保する上で不可欠な元素であり、その下限は１．７％である。しかし、Ｍｎが多すぎると鋼の焼入れ性が増してＨＡＺ靱性、現地溶接性を劣化させるだけでなく、連続鋳造鋼片の中心偏析を助長し、母材の低温靱性をも劣化させるので上限を２．５％とした。
【００１２】
本発明では、不純物元素であるＰ、Ｓ量をそれぞれ０．０１５％、０．００３％以下とする。この主たる理由は母材およびＨＡＺの低温靱性をより一層向上させるためである。Ｐ量の低減は連続鋳造スラブの中心偏析を軽減するとともに、粒界破壊を防止して低温靱性を向上させる。また、Ｓ量の低減は熱間圧延で延伸化するＭｎＳを低減して延靱性を向上させる効果がある。
【００１３】
Ｎｉを添加する目的は低炭素の本発明鋼を低温靱性や現地溶接性を劣化させることなく向上させるためである。Ｎｉ添加はＭｎやＣｒ、Ｍｏ添加に比較して圧延組織（とくに連続鋳造鋼片の中心偏析帯）中に低温靱性に有害な硬化組織を形成することが少ないばかりか、０．１％以上の微量Ｎｉ添加がＨＡＺ靱性の改善にも有効であることが判明した（ＨＡＺ靱性上、とくに有効なＮｉ添加量は０．３％以上である）。しかし、添加量が多すぎると、経済性だけでなく、ＨＡＺ靱性や現地溶接性を劣化させるので、その上限を１．０％とした。また、Ｎｉ添加は連続鋳造時、熱間圧延時におけるＣｕ割れの防止にも有効である。この場合、ＮｉはＣｕ量の１／３以上添加する必要がある。
【００１４】
Ｍｏを添加する理由は鋼の焼入れ性を向上させ、目的とするマルテンサイト主体の組織を得るためである。Ｂ添加鋼においてはＭｏの焼入れ性向上効果が高まり、また、ＭｏはＮｂと共存して制御圧延時にオーステナイトの再結晶を抑制し、オーステナイト組織の微細化にも効果がある。このような効果を得るために、Ｍｏは最低でも０．１５％必要である。しかし、過剰なＭｏ添加はＨＡＺ靱性、現地溶接性を劣化させ、さらにＢの焼入れ性向上効果を消失せしめることもあるので、その上限を０．６％とした。
【００１５】
Ｎｂは、０．０１〜０．１０％を含有する。ＮｂはＭｏと共存して制御圧延時にオーステナイトの再結晶を抑制して組織を微細化するだけでなく、析出硬化や焼入れ性増大にも寄与し、鋼を強靱化するため、０．０１％以上含有する。特にＮｂとＢが共存すると焼入れ性向上効果が相乗的に高まる。しかし、Ｎｂ添加量が多すぎると、ＨＡＺ靱性や現地溶接性に悪影響をもたらすので、その上限を０．１０％とした。
【００１６】
Ｔｉは、０．００５〜０．０３０％を含有する。Ｔｉ添加は微細なＴｉＮを形成し、スラブ再加熱時およびＨＡＺのオーステナイト粒の粗大化を抑制してミクロ組織を微細化し、母材およびＨＡＺの低温靱性を改善する。また、Ｂの焼入れ性向上効果に有害な固溶ＮをＴｉＮとして固定する役割も有する。この目的のために、Ｔｉ量は３．４Ｎ（各々重量％）以上添加することが望ましい。また、Ａｌ量が少ない時（たとえば０．００５％以下）、Ｔｉは酸化物を形成し、ＨＡＺにおいて粒内フェライト生成核として作用し、ＨＡＺ組織を微細化する効果も有する。このようなＴｉＮの効果を発現させるためには、最低０．００５％のＴｉ添加が必要である。しかし、Ｔｉ量が多すぎると、ＴｉＮの粗大化やＴｉＣによる析出硬化が生じ、低温靱性を劣化させるので、その上限を０．０３０％に限定した。
【００１７】
Ａｌは通常脱酸材として鋼に含まれる元素で、組織の微細化にも効果を有する。しかし、Ａｌ量が０．０６％を越えるとＡｌ系非金属介在物が増加して鋼の清浄度を害するので、上限を０．０６％とした。しかし、脱酸はＴｉあるいはＳｉでも可能であり、Ａｌは必ずしも添加する必要はない。
以上は、本発明の鋼管母材の主要成分であるが、必要に応じて以下の成分を選択的に含有させる。
【００１８】
Ｂは極微量で鋼の焼入れ性を飛躍的に高め、目的とするマルテンサイト主体の組織を得るために、非常に有効な元素である。さらに、ＢはＭｏの焼入れ性向上効果を高めると共に、Ｎｂと共存して相乗的に焼入れ性を増す。一方、過剰に添加すると、低温靱性を劣化させるだけでなく、かえってＢの焼入れ性向上効果を消失せしめることもあるので、その上限を０．００３０％とした。
【００１９】
ＮはＴｉＮを形成しスラブ再加熱時およびＨＡＺのオーステナイト粒の粗大化を抑制して母材、ＨＡＺの低温靱性を向上させる。このために必要な最小量は０．００１％である。しかし、Ｎ量が多すぎるとスラブ表面疵や固溶ＮによるＨＡＺ靱性の劣化、Ｂの焼入れ性向上効果の低下の原因となるので、その上限は０．００６％に抑える必要がある。
【００２０】
つぎに、Ｖ、Ｃｕ、Ｃｒ、Ｃａ、 ＲＥＭ、 Ｍｇを添加する目的について説明する。
本発明の鋼管母材の基本成分に、更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・靱性の一層の向上や製造可能な鋼材サイズの拡大をはかるためである。したがって、その添加量は自ずから制限されるべき性質のものである。
【００２１】
ＶはＮｂとほぼ同様の効果を有するが、その効果はＮｂに比較して弱い。しかし、超高強度鋼におけるＶ添加の効果は大きく、ＮｂとＶの複合添加は本発明鋼の優れた特徴をさらに顕著なものとする。上限はＨＡＺ靱性、現地溶接性の点から０．１０％まで許容できるが、特に０．０３〜０．０８％の添加が望ましい範囲である。
【００２２】
Ｃｕは母材、溶接部の強度を増加させるが、多すぎるとＨＡＺ靱性や現地溶接性を著しく劣化させる。このためＣｕ量の上限は１．０％である。
Ｃｒは母材、溶接部の強度を増加させるが、多すぎるとＨＡＺ靱性や現地溶接性を著しく劣化させる。このためＣｒ量の上限は０．６％である。
ＣａおよびＲＥＭは硫化物（ＭｎＳ）の形態を制御し、低温靱性を向上（シャルピー試験の吸収エネルギーの増加など）させる。Ｃａ量が０．００６％、ＲＥＭが０．０２％を越えて添加するとＣａＯ−ＣａＳまたはＲＥＭ−ＣａＳが大量に生成して大型クラスター、大型介在物となり、鋼の清浄度を害するだけでなく、現地溶接性にも悪影響をおよぼす。このためＣａ添加量の上限を０．００６％またはＲＥＭ添加量の条件を０．０２％に制限した。なお超高強度ラインパイプでは、Ｓ、Ｏ量をそれぞれ０．００１％、０．００２％以下に低減し、以下に示すＭｎＳのクラスターの形状を制御するための指標であるＥＳＳＰ（Effestive Sulphide Shape Controlling Parameter）が０．５≦ＥＳＳＰ≦１０．０を満足するようにＣａ、Ｓ、Ｏを調整することがとくに有効である。
【００２３】
ＥＳＳＰ＝（Ｃａ）〔１−１２４（Ｏ）〕／１．２５Ｓ … （１）
上記のＥＳＳＰが０．５未満になるとＣａＯ−ＣａＳが大量の生成して粗大なクラスター、粗大介在物となり溶接割れ等の溶接性を悪化させ、上記ＥＳＳＰが１０．０を越えると、ＭｎＳの形状制御の効果がなくなるため、ＥＳＳＰを０．５〜１０．０に規定する。
【００２４】
Ｍｇは微細分散した酸化物を形成し、溶接熱影響部の粒粗大化を抑制して低温靭性を向上させる。０．００６％以上では粗大酸化物を生成し逆に靭性を劣化させる。
以上の個々の添加元素の限定に加えて、さらに以下に示す焼き入れ性を表す指標であるＰを１．９≦Ｐ≦４．０に制限することが望ましい。
【００２５】
Ｐ＝２．７Ｃ＋０．４Ｓｉ＋Ｍｎ＋０．８Ｃｒ＋０．４５（Ｎｉ＋Ｃｕ）＋（１＋β）Ｍｏ−１＋β … （２）
但し、Ｂ≧３ｐｐｍではβ＝１、Ｂ＜３ｐｐｍではβ＝０とする。
Ｐを上記のように制御する理由は、目的とする強度・低温靱性バランスを達成するためである。Ｐ値の下限を１．９としたのは９００ＭＰａ以上の強度と優れた低温靱性を得るためである。また、Ｐ値の上限を４．０としたのは優れたＨＡＺ靱性、現地溶接性を維持するためである。
【００２６】
以上が本発明の鋼管母材に含有する成分の限定根拠であるが、以上のような化学成分を有していても、本発明の組織である微細なマルテンサイト＋ベイナイト主体の組織が得られるための適正な製造条件としなければ所望の特性は得られない。微細なマルテンサイト主体の組織を得る原理的な方法は、再結晶粒を未再結晶温度域で加工し、板厚方向に偏平したオーステナイト粒とし、これをフェライト生成が抑制される臨界冷却速度以上の冷却速度で冷却することである。
【００２７】
望ましい製造方法は、本発明の化学成分を有する鋼片を９５０〜１２５０℃に再加熱し、７００〜９５０℃での累積圧下量が５０％以上となるように７００℃以上の鋼材温度で圧延した後、１０℃以上の冷却速度で５５０℃以下まで冷却する。 また必要に応じてＡC1変態点以下の温度で焼戻しを行う。
本発明の鋼管は、このようにして製造された鋼板を管状に成形した後、鋼板両端部の突き合わせ部をアーク溶接し、さらに拡管して鋼管をする。
【００２８】
次ぎに、本発明の鋼管の溶接金属部の成分の限定理由について述べる。
Ｃ量は、０．０３〜０．１４％に限定する。炭素は鋼の強度向上に極めて有効であり、マルテンサイト組織において目標とする強度を得るためには、最低０．０３％は必要である。しかし、Ｃ量が多すぎると溶接低温割れが発生しやすくなり、現地溶接部とシーム溶接が交わるいわゆるＴクロス部のＨＡＺの最高硬さの上昇招くので、その上限を０．１４％とした。さらに、望ましくは上限値は０．１０％が好ましい。
【００２９】
Ｓｉはブローホール防止のために０．０５％以上は必要であるが、含有量が多いと低温靱性を著しく劣化させるので、上限を０．４０％とした。特に、内外面溶接や多層溶接を行う場合、再熱部の低温靱性を劣化させる。
Ｍｎは優れた強度・低温靱性のバランスを確保する上で不可欠な元素であり、その下限は１．２％である。しかし、Ｍｎが多すぎると偏析が助長され低温靱性を劣化させるだけでなく、溶接材料の製造も困難になるので上限を２．２％とした。
【００３０】
Ｐ、Ｓは、低温靭性の劣化、低温割れ感受性の低減のために、Ｐ、Ｓの量は低い方が望ましく、上限量をそれぞれ０．０１０％と規定した。
Ｎｉを添加する目的は焼入れ性を高めて強度を確保し、さらに低温靱性向上させるためである。１．３％以下では目標の強度、低温靭性を得ることが難しい。一方、含有量が多すぎると高温割れの危険があるため上限は３．２％とした。
【００３１】
Ｃｒ、Ｍｏ、Ｖの効果の違いは厳密には区別できないが、いずれも焼入れ性を高めることにより高強度を得るために添加する。Ｃｒ、Ｍｏ、Ｖの１種又は２種以上の合計量が１．０％以下では効果が十分でなく、一方多量に添加すると低温割れの危険が増すため上限を２．５％とした。
Ｂは微量で焼入れ性を高め、溶接金属の低温靭性向上に有効な元素であるが、含有量が多すぎると却って低温靭性が低下するので含有範囲を０．００５％以下とした。
【００３２】
溶接金属には、その他の成分として、溶接時の精錬・凝固を良好に行わせるために必要に応じて添加されたＴｉ, Ａｌ，Ｚｒ，Ｎｂ，Ｍｇ等の元素を含有する場合があるが、残部は鉄および不可避的不純物である。
本発明の超高強度鋼板は、先に述べた成分を規定した鋼を鋳造後、熱間加工し、その後急冷したり、場合によっては焼戻しを行って製造される。引張強さ９００ＭＰａ以上の超高強度を達成するためには、鋼をマルテンサイト・ベイナイト等の低温変態組織主体のミクロ組織にしてフェライトの生成を抑制する必要がある。
【００３３】
溶接金属は、溶接後の凝固まま組織であり、冷却速度が遅い溶接金属において、本発明が上記の目的強度を得、さらに本発明の鋼板と同様に優れた低温靱性を得るためには、溶接金属の化学成分と組織の調整が必要である。
Ｎｉは焼入性を高めて低い冷却速度でも高強度を得ることを可能にし、また、マルテンサイトラス間に残留オーステナイトを形成することを促進し低温靱性を向上させる。
【００３４】
本発明では、溶接金属のＮｉ量を鋼板成分より１％以上高くし、かつ、溶接金属部ならびに溶接熱影響部をベイナイト・マルテンサイト組織にすることにより、所望の強度と低温靱性が得られる。溶接金属のＮｉ量が鋼板成分より１％低い場合は、上記効果が得られないため、本発明では、その下限を１％とした。
次に、本発明の鋼管を製造する方法について説明する。
【００３５】
本願発明が目指すラインパイプは通常、直径が４５０ｍｍから１５００ｍｍ、肉厚が１０ｍｍから４０ｍｍ程度のサイズである。このようなサイズの鋼管を高率良く製造する方法としては、鋼板をＵ形次いでＯ形に成形するＵＯ工程で製管し、鋼板の両端部を突き合わせて、突き合わせ部をＭＡＧアーク溶接等で外面から仮付け溶接した後に、この仮付け溶接部を内外面からサブマージアーク溶接等で本溶接し、その後、拡管して真円度を高める鋼管の製造方法が確立されている。
【００３６】
サブマージアーク溶接等のアーク溶接方法は母材の希釈が大きい溶接であり、所望の特性すなわち溶接金属組成を得るためには、母材の希釈を考慮した溶接材料の選択が必要である。
以下に、本発明の超高強度ラインパイプを製造する際の溶接に用いる溶接ワイヤーの化学組成の限定理由を述べる。なお、本発明の本溶接の前に行う鋼板付き合わせ部の仮付け溶接は、溶接面積が少なく本溶接に比べて溶接金属部の品質の影響が小さいため、本溶接に用いる溶接ワイヤーの成分は、すべて以下のように規定するが、仮付け溶接に用いる溶接ワイヤーの成分はＣ、Ｓｉ、Ｍｎ以外の他の成分は、特に規定する必要はない。
【００３７】
Ｃは、溶接金属で必要とされるＣ量の範囲を得るために、母材成分による希釈および雰囲気からＣの混入を考慮して０．０１〜０．１２％とした。
Ｓｉは、溶接金属で必要とされるＳｉ量の範囲を得るために、母材成分による希釈を考慮して０．３％以下とした。
Ｍｎは、溶接金属で必要とされるＭｎ量の範囲を得るために、母材成分による希釈を考慮して１．２％〜２．４％とした。
【００３８】
Ｎｉは、溶接金属で必要とされるＮｉ量の範囲を得るために、母材成分による希釈を考慮して４．０％〜８．５％とした。
Ｃｒ、Ｍｏ、Ｖは、溶接金属で必要とされるＣｒ、Ｍｏ、Ｖのうちの１種又は２種以上の合計量の範囲を得るために、母材成分による希釈を考慮して３．０％〜５．０％とした。
【００３９】
その他Ｐ，Ｓの不純物は極力少ない方が望ましく、Ｂは強度確保に添加することも可能である。また、Ｔｉ，Ａｌ，Ｚｒ，Ｎｂ，Ｍｇ等が脱酸を目的として使用される。
なお、本発明の仮付け溶接及び本溶接は、単極だけでなく、複数電極での溶接も可能である。複数電極で溶接の場合は各種ワイヤーの組み合わせが可能であり、個々のワイヤーが上記成分範囲にある必要はなく、それぞれのワイヤー成分と消費量からの平均組成が上記成分範囲にあれば良い。
【００４０】
サブマージアーク溶接等の本溶接に使用されるフラックスは大別すると焼成型フラックスと溶融型フラックスがある。焼成型フラックスは合金材添加が可能で拡散性水素量が低い利点があるが、粉化しやすく繰り返し使用が難しい欠点がある。一方、溶融型フラックスはガラス粉状で、粒強度が高く、吸湿しにくい利点があり、拡散性水素がやや高い欠点がある。本願発明のごとき超高強度の場合は、溶接低温割れが起こりやすく、この点からは焼成型が望ましいが、一方、回収して繰り返し使用が可能な溶融型は大量生産に向きコストが低い利点がある。焼成型ではコストが高いことが、溶融型では厳密な品質管理の必要性が問題であるが、工業的に対処可能な範囲であり、どちらでも本質的には使用可能である。
【００４１】
溶接条件については望ましい範囲は以下の通りである。
最初に行う仮付け溶接は、ＭＡＧアーク溶接でもＭＩＧアーク溶接でもＴＩＧアーク溶接でもよい。通常はＭＡＧアーク溶接である。次に仮付け溶接後に行う本溶接は、通常サブマージドアーク溶接であるが、ＴＩＧアーク溶接でもＭＩＧアーク溶接でも、ＭＡＧアーク溶接でもよい。溶接速度は１〜３ｍ／分程度が適切な範囲である。１ｍ／分未満の溶接はラインパイプのシーム溶接としては非効率であり、３ｍ／分を超える高速溶接ではビード形状が安定しない。仮付け溶接とその後の本溶接が重複するならば溶接入熱は出来る限り低い方が好ましい。また、本溶接のアーク溶接は何パスでも行ってもよい。溶接入熱は板厚によって異なるが、例えば板厚１６ｍｍの場合では溶接入熱を１．０〜２．７ｋＪ／ｍｍにすることが望ましい範囲である。入熱が小さすぎると溶け込みが不十分になり、溶接回数が多くなり、作業効率が悪くなり、溶接入熱が大きすぎると熱影響部の軟化が大きく、溶接部の靭性も低下する。
【００４２】
これらのシーム溶接後、拡管により真円度を向上させる。真円にするためには塑性域まで変形させる必要があるが、本願発明のごとき高強度鋼の場合は０．７％程度以上の拡管率（＝（拡管後円周−拡管前円周）／拡管前円周）が必要であるが、２％を超える大きな拡管を行うと、母材、溶接部とも塑性変形による靭性劣化が大きくなるため、拡管率は０．７〜２％以下にするのが望ましい。
【００４３】
【実施例】
以下に、本発明の実施例とその効果を具体的に説明する。
表１に示す本発明範囲を満たす成分の発明鋼（Ａ鋼〜Ｄ鋼）及び本発明範囲を外れる成分の比較鋼（Ｅ鋼，Ｆ鋼）を３００トン転炉で溶製後、連続鋳造鋼片とし、その後１１００℃に再加熱後、再結晶域で圧延し、その後９００〜７５０℃の累積圧下量が７５％となる制御圧延を１６ｍｍまで行い、その後水冷停止温度が２００〜４５０℃になるように水冷して鋼板を製造した。その結果、表１に示されるように発明鋼（Ａ鋼〜Ｄ鋼）の鋼板の強度は、本発明の目標範囲（９００ＭＰａ以上）となり、低温靭性（シャルピー試験の−３０℃での吸収エネルギー：２３０Ｊ以上）も高かった。一方、Ｃ量が高くＮｉが添加されていないＥ鋼の鋼板の強度は、本発明の目標範囲にあるが、低温靭性が低くなり、Ｃ量が低いＦ鋼の鋼板の低温靱性は目標範囲になるが、強度が低い。
【００４４】
【表１】
【表２】
これらの鋼板をさらにＵＯ工場で管状に成形し、鋼板の付き合わせ部を８０％Ａｒ＋２０％ＣＯ2 のシールドガスでMAG アーク溶接を用いて仮付け溶接を行った後、表２に示す溶接ワイヤー及びフラックスを用い３電極、２．０ｍ／分、入熱１．５ＫＪ／ｍｍの溶接条件で仮溶接部の内外面を各１パスのサブマージアーク溶接による本溶接を行い、その後、拡管率１％の拡管を行った。得られた鋼管の特性を評価した結果を表３に示す。
【００４７】
表２に示す本発明範囲を満たす成分の鋼及び溶接ワイヤーを用いて溶接した発明例（実施Ｎｏ．１〜６）では、鋼管シーム溶接部に良好な溶接ビードが得られ、溶接金属部の化学成分は本発明の範囲を満たし、溶接金属強度（９００ＭＰａ以上）、溶接金属の引張強度と鋼板の引張強度の差も適正範囲（−１００ＭＰａ以上）も適性であり、本溶接によって形成される内面溶接金属部と外面溶接金属部の間隔（Δｄ）（Δｄ：０ｍｍ超））も適性であった。また、これらの本発明範囲を満たす発明例の鋼管は、母材部及び溶接部が共に目標とする強度、低温靱性等の機械的性質を有し、バースト試験においても管体破断が達成できた。
【００４８】
一方、比較例の実施Ｎｏ．７〜９は、母材成分は本発明の範囲であるが、ワイヤー成分が本発明の範囲外（Ｎｏ．７：Ｎｉが低目、Ｎｏ．８：Ｃが高目、Ｎｏ．９：Ｎｉが高目）であるため、溶接金属部の成分が本発明範囲を外れた。その結果、Ｎｏ．７は溶接金属の強度が低くなり、また、溶接金属の引張強度と鋼板の引張強度の差が適正範囲（−１００ＭＰａ以上）を外れたためバースト試験では溶接部破断が生じた。また、Ｎｏ．８では溶接部の低温割れが発生し、Ｎｏ．９は高温割れが発生したため、引張り試験、バースト試験は実施できなかった。
【００４９】
比較例の実施Ｎｏ．１０は、溶接ワイヤーの成分は本発明の範囲内であるが、鋼板の成分が本発明範囲外であるため、鋼管母材の低温靱性が目標（低目）に達しなかった。
しかしながら、比較例のＮｏ．１１及びＮｏ．１２では、母材及び溶接ワイヤーの成分は本発明の範囲内であるが、本溶接で形成された溶接部中の内面溶接金属部と外面溶接金属部との間隔（Δｄ）が本発明の範囲（Δｄ＞０）を外れたためバースト試験では溶接熱影響部破断が生じた。また、比較例のＮｏ．１３は、溶接材料及び溶接金属のＣが本願発明範囲を外れている（高目）ために、溶接金属の靭性が低いためにラインパイプの要求特性を満たしていない。比較例のＮｏ１４は母材及び溶接ワイヤーの成分は本発明の範囲内であるが、溶接金属の引張強度と鋼板の引張強度の差が適正範囲（−１００ＭＰａ以上）を外れたためバースト試験では溶接熱影響部破断が生じた。
【００５０】
【表３】
【表４】
【表５】
【発明の効果】
本発明によれば、母材、溶接部共に低温靱性のバランスが優れ、かつ現地溶接が容易な引張強さ９００ＭＰａ以上（ＡＰＩ規格Ｘ１００超）の超高強度ラインパイプが実現可能であり、長距離パイプラインの敷設コストが低下し、世界のエネルギー問題解決に寄与できる。
【図面の簡単な説明】
【図１】本発明の鋼管シーム溶接部の断面図。
【図２】従来の鋼管シーム溶接部の断面図。
【符号の説明】
１…本溶接金属部の外面溶接金属部
２…本溶接金属部の内面溶接金属部
３…鋼管の母材部
４…仮付け溶接金属部[0001]
BACKGROUND OF THE INVENTION
The seam having a tensile strength of 900 MPa or more, which can be widely used as a line pipe for transporting natural gas and crude oil, can improve transportation efficiency by increasing pressure, and can improve local construction efficiency by reducing outer diameter and weight. The present invention relates to an ultra-high-strength line pipe excellent in low-temperature toughness of a welded portion and a manufacturing method thereof.
[0002]
[Prior art]
In recent years, pipelines have become increasingly important as long-distance transportation methods for crude oil and natural gas. Currently, the American Petroleum Institute (API) standard X65 is the basic design for trunk line pipes for long-distance transportation, and the actual usage is overwhelmingly large. However, higher strength line pipes are required for (1) improving transportation efficiency by increasing pressure, and (2) improving local construction efficiency by reducing the outer diameter and weight of the line pipe. Up to now, line pipes up to X80 (tensile strength of 620 MPa or more) have been put into practical use, but the need for higher-strength line pipes has become stronger. Currently, research on ultra-high-strength line pipe manufacturing methods is based on conventional X80 line pipe manufacturing techniques (eg, NKK Technical Report No. 138 (1992), pp24-31 and The 7th Offshore Mechanics and Arctic Engineering (1988), Volume V, pp 179-185), but it is considered that the production of X100 (tensile strength of 760 MPa or more) line pipe is the limit. For ultra-high-strength line pipes exceeding X100, steel plate manufacturing research has already been conducted (PCT / JP96 / 00155, 00157). However, conventional technology relating to seam welding cannot be applied to such an ultra-high-strength line pipe, and if the problems relating to seam welding technology cannot be solved, it is impossible to produce a steel pipe even if a steel plate can be produced. Ultra-high strength pipelines have many problems such as welded heat affected zone (HAZ) toughness, on-site weldability, joint softening, and tube breakage due to burst tests, including the balance between strength and low temperature toughness. There is a demand for early development of a revolutionary ultra-high-strength line pipe (exceeding X100).
[0003]
[Problems to be solved by the invention]
The present invention has an excellent balance of low-temperature toughness and is easy to be welded on site, and has a tensile strength (TS) of 900 MPa or more (API standard X100). It is an object of the present invention to provide an ultra-high-strength line pipe excellent in low-temperature toughness of a seam weld and a manufacturing method thereof.
[0004]
[Means for Solving the Problems]
The gist of the present invention is as follows.
(1) A seam welded steel pipe in which the tensile strength of the base metal part is 900 MPa or more and the difference between the tensile strength of the weld metal part and the tensile strength of the base material part is -100 MPa or more, In the weld metal portion, the interval between the inner surface weld metal portion and the outer surface weld metal portion formed by the main welding performed after the tack welding of the steel plate mating portion in the pipe making process is greater than 0 mm, and the inner surface weld metal portion A super-high-strength welded steel pipe excellent in low-temperature toughness of the seam welded portion, wherein the welded metal portion overlapped with the tack welded metal portion formed by the tack welding.
(2) By weight%
C: 0.03-0.10%,
Si: 0.6% or less,
Mn: 1.7-2.5%,
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 1.0%,
Mo: 0.15-0.60%,
Nb: 0.01-0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less,
In addition, by weight%, B: 0.0030% or less, N: 0.001 to 0.006%, V: 0.10% or less, Cu: 1.0% or less, Cr: 1.0% Hereinafter, Ca: 0.01% or less, REM: 0.02% or less, Mg: 0.006% or less of one or two or more of the base material part consisting of iron and inevitable impurities,
% By weight
C: 0.03-0.14%,
Si: 0.05 to 040%,
Mn: 1.2-2.2%,
P: 0.010% or less,
S: 0.010% or less,
Ni: 1.3-3.2%
The total amount of one or more of Cr, Mo and V is 1.0 to 2.5%,
B: 0.005% or less,
And the balance consists of a weld metal part consisting of iron and inevitable impurities, and the amount of Ni in the weld metal part is 1% or more higher than the amount of Ni in the base metal part. The super high strength welded steel pipe excellent in low temperature toughness of the seam welded portion according to (1) above, wherein the structure of the seam welded portion including the weld heat affected zone is composed of bainite martensite.
(3) By weight%
C: 0.03-0.10%,
Si: 0.6% or less,
Mn: 1.7-2.5%,
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 1.0%,
Mo: 0.15-0.60%,
Nb: 0.01-0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less,
In addition, by weight%, B: 0.0030% or less, N: 0.001 to 0.006%, V: 0.10% or less, Cu: 1.0% or less, Cr: 1.0% Hereinafter, both ends of a steel sheet containing one or more of Ca: 0.01% or less, REM: 0.02% or less, Mg: 0.006% or less, the balance being iron and unavoidable impurities. After the attachment, the attachment part contains C: 0.01 to 0.12%, Si: 0.3% or less, and Mn: 1.2 to 2.4% by weight%. After performing tack welding from the outer surface using a welding wire as a component, the tack welded portion is expressed in terms of% by weight, C: 0.01 to 0.12%, Si: 0.3% or less, Mn : 1.2 to 2.4%, Ni: 4.0 to 8.5%, containing one or more of Cr, Mo, V or a total amount of 3.0 to 5.0%, or The distance between the inner weld metal part and the outer weld metal part formed by welding is greater than 0 mm using a welding wire and a flux mainly composed of Fe whose Ni content is 1% or more higher than the Ni content of the steel sheet. And performing the main welding from the inner surface and the outer surface so that the inner weld metal portion and the outer weld metal portion overlap with the tack weld metal portion formed by the tack welding, respectively. A method for producing an ultra-high-strength welded steel pipe excellent in low temperature toughness of a seam weld.
(4) In the main welding, the tack welded portion is welded with two or more passes from the inner surface and the outer surface, respectively. Steel pipe manufacturing method.
(5) As the tack welding, any one of MAG arc welding, MIG arc welding, and TIG arc welding is used, and as the main welding, submerged arc welding, MAG arc welding, MIG arc welding, TIG arc welding are used. Any one method is used, The manufacturing method of the ultra high strength welded steel pipe excellent in the low temperature toughness of the seam welded part according to any one of the above (3) or (4).
[0005]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the contents of the present invention will be described in detail.
The present invention relates to an ultra-high-strength line pipe excellent in low-temperature toughness of a seam weld having a tensile strength (TS) of 900 MPa or more. The ultra-high-strength line pipe of this strength level can withstand about twice as much pressure as the conventional mainstream X65, and therefore can transport about twice as much gas at the same size. On the other hand, when using X65 to achieve the same gas transport efficiency as the above ultra-high-strength line pipe, it is necessary to increase the wall thickness in order to increase the pressure, resulting in higher material costs, transport costs, and local welding costs. Pipeline laying costs will increase significantly. This is the reason why an ultra-high-strength line pipe excellent in low-temperature toughness having a tensile strength (TS) of 900 MPa or more is required. Conventionally, in such an ultra-high-strength line pipe having a tensile strength of 900 MPa or more, it is extremely difficult to manufacture a steel pipe, and it is difficult to ensure the characteristics of the steel pipe particularly at low temperature toughness. As a guideline to guarantee the target characteristics including the seam welded part of the steel pipe, in the burst test, the fracture at the pipe body is achieved without breaking at the weld heat affected zone and weld metal, etc. Improving toughness is an important technical challenge. In conventional ultra-high-strength line pipes, when welding, coarse MA (Martensite-Austenite Constituent: Martensite-Austenite Constituent) along the former austenite grain boundary from the meeting part of the heat-affected zone of the seam weld to 1 mm This is a factor that significantly reduces the absorbed energy value. Therefore, the V-notch Charpy absorbed energy at the meeting part or meeting part + 1 mm of the conventional welding heat-affected zone is as low as less than 50 J at −30 ° C., for example, to satisfy the target of 64 J or more at −30 ° C. Was quite difficult.
[0006]
The present inventors diligently studied through experiments and the like in order to improve the low temperature toughness of the seam welded part in an ultrahigh strength line pipe having a tensile strength of 900 MPa or more.
1 and 2 show a conventional welded portion and a welded portion according to the present invention in an ultra-high-strength line pipe. Usually, seam welding at the time of pipe making of steel pipes involves attaching both ends of a steel plate, and then temporarily attaching the attached portion from the outer surface by MAG arc welding or the like (hereinafter referred to as “tack welding”). The welded portion is further welded from the inner surface and the outer surface by submerged arc welding or the like (hereinafter referred to as main welding). At this time, as shown in FIG. 1, conventional main welding includes a weld metal portion formed by welding from the outer surface in the main welding (hereinafter referred to as an outer surface weld metal portion) and a weld metal portion formed by welding from the inner surface ( (Hereinafter referred to as “inner weld metal part”), the tack weld metal part made before the main welding melts and disappears, the welding heat input becomes excessively high, and the old austenite grains in the weld heat affected zone become coarse. At the same time, the MA produced along the prior austenite grain boundaries also becomes coarse, which reduces the Charpy absorbed energy of the weld heat affected zone, and softens the weld heat affected zone, causing breakage from the welded zone in the burst test (pipe). It turned out to be a factor incurring (not body rupture).
[0007]
The present inventors have a problem that the low temperature toughness of the welded portion due to an excessive increase in welding heat input in the main welding from the inner and outer surfaces where the outer surface weld metal portion and the inner surface weld metal portion overlap each other as in the prior art. In order to solve this problem, detailed examination was made on the optimum conditions for the main welding. As a result, as shown in FIG. 2, in the main welding after the tack welding, the outer welding metal part and the inner welding metal part are not overlapped, and the temporary welding metal part is left without being melted. By avoiding a significant increase in welding heat input, coarsening of the prior austenite grain size and MA coarsening in the coarse-grained part in the welded heat-affected zone is suppressed. It was found that the tube notch breakage (no breakage from the welded portion) can be achieved in the burst test by improving the V-notch Charpy absorbed energy in and by suppressing the softened portion of the weld heat affected zone. Also, when performing such main welding, in order to prevent the occurrence of weld defects in the welded portion, each of the inner surface welded metal portion and the outer surface welded metal portion formed by the main welding and the temporary welding before that It was found that it was necessary to overlap the formed temporary weld metal part.
[0008]
From the above knowledge, in the present invention, in the seam welded portion of a line pipe with an ultrahigh strength having a tensile strength of 900 MPa or more, the distance between the inner surface welded metal portion and the outer surface welded metal portion formed by the main welding after the tack welding. (Δd) is greater than 0 mm (in contrast, the conventional Δd is a value of 0 mm or less), and the inner surface weld metal portion and the outer surface weld metal portion are formed by the tack welding. It is a requirement to overlap each other. When Δd is 0 mm or less, that is, when the inner surface weld metal part and the outer surface weld metal part formed by the main welding overlap, the temporary weld metal part performed before the main welding as described above melts and disappears, As a result, the heat becomes excessively high, and as a result, the old austenite grains in the weld heat affected zone become coarser, and the MA formed along the former austenite grain boundary also becomes coarser, resulting in a decrease in Charpy absorbed energy and welding heat in the weld heat affected zone. The affected area softens. Moreover, if each of the inner surface weld metal part and outer surface weld metal part formed by this welding part and the tack welding metal part formed by tack welding do not overlap, the welding defect of a welding part will generate | occur | produce. . Here, the method of the tack welding first performed in the present invention may be a commonly known MAG arc welding, MIG arc welding, or TIG welding. Also, inner / outer surface welding performed after tack welding may be submerged arc welding, MIG arc welding, or TIG welding. Further, the welding from the inner surface and the outer surface of the tack welded portion in the main welding after the tack welding may be either one-pass welding or two-pass welding, but the welding heat input during the main welding It is more preferable to perform welding with two or more passes from the viewpoint of suppressing the softening of the heat affected zone by lowering as much as possible.
[0009]
Further, from the results of numerous burst tests using the steel pipe manufactured by the above-described welding method, the inventors of the present invention have a tensile strength of the weld metal of [base material strength] −100 (MPa) or more. It has been found that it does not break from the weld, but breaks from the tube. Accordingly, in the present invention, the average tensile strength of the welded portion is set to [circumferential tensile strength in the base material portion] −100 (MPa) or more.
[0010]
Next, components and structures of the base metal part and the weld metal part constituting the steel pipe of the present invention will be described.
First, the reasons for limiting the base material components of the present invention are as follows.
The amount of C is limited to 0.03 to 0.10%. Carbon is extremely effective in improving the strength of steel, and at least 0.03% is necessary to obtain the target strength in the martensite structure. However, if the amount of C is too large, the base material, HAZ low temperature toughness and on-site weldability are significantly deteriorated, so the upper limit was made 0.10%. Furthermore, the upper limit is preferably 0.08%.
[0011]
Si is an element added for deoxidation and strength improvement, but if added in a large amount, the HAZ toughness and on-site weldability are remarkably deteriorated, so the upper limit was made 0.6%. Steel can be deoxidized with either Al or Ti, and Si does not necessarily have to be added.
Mn is a martensite-based microstructure of the steel of the present invention and is an indispensable element for securing an excellent balance between strength and low temperature toughness, and its lower limit is 1.7%. However, if Mn is too much, not only the hardenability of the steel will increase and the HAZ toughness and on-site weldability will deteriorate, but also the center segregation of the continuously cast steel slab will be promoted, and the low temperature toughness of the base metal will also be deteriorated. 2.5%.
[0012]
In the present invention, the amounts of impurity elements P and S are set to 0.015% and 0.003% or less, respectively. The main reason for this is to further improve the low temperature toughness of the base material and the HAZ. The reduction of the amount of P reduces the center segregation of the continuously cast slab and prevents the grain boundary fracture, thereby improving the low temperature toughness. Further, the reduction of the amount of S has the effect of reducing the MnS stretched by hot rolling and improving the ductility.
[0013]
The purpose of adding Ni is to improve the low carbon steel of the present invention without deteriorating the low temperature toughness and on-site weldability. Compared with the addition of Mn, Cr and Mo, the addition of Ni is less likely to form a hardened structure harmful to low temperature toughness in the rolled structure (especially the central segregation zone of the continuous cast steel slab). It has been found that the addition of a trace amount of Ni is also effective for improving the HAZ toughness (in particular, the effective Ni addition amount is 0.3% or more in terms of HAZ toughness). However, if the addition amount is too large, not only the economy but also the HAZ toughness and on-site weldability are deteriorated, so the upper limit was made 1.0%. Ni addition is also effective for preventing Cu cracking during continuous casting and hot rolling. In this case, Ni needs to be added by 1/3 or more of the amount of Cu.
[0014]
The reason for adding Mo is to improve the hardenability of the steel and to obtain the target martensite-based structure. In the B-added steel, the effect of improving the hardenability of Mo is enhanced, and Mo coexists with Nb to suppress recrystallization of austenite during controlled rolling, and is effective in refining the austenite structure. In order to obtain such an effect, Mo needs to be at least 0.15%. However, excessive addition of Mo deteriorates the HAZ toughness and on-site weldability, and may further eliminate the effect of improving the hardenability of B, so the upper limit was made 0.6%.
[0015]
Nb contains 0.01 to 0.10%. Nb coexists with Mo to suppress the recrystallization of austenite during controlled rolling to refine the structure, contribute to precipitation hardening and hardenability, and toughen the steel. contains. In particular, when Nb and B coexist, the effect of improving hardenability increases synergistically. However, if the amount of Nb added is too large, the HAZ toughness and on-site weldability are adversely affected, so the upper limit was made 0.10%.
[0016]
Ti contains 0.005-0.030%. Addition of Ti forms fine TiN, suppresses coarsening of austenite grains during slab reheating and HAZ, refines the microstructure, and improves the low temperature toughness of the base material and HAZ. Moreover, it has a role which fixes solid solution N harmful to the hardenability improvement effect of B as TiN. For this purpose, it is desirable to add Ti in an amount of 3.4 N (each by weight%) or more. Further, when the amount of Al is small (for example, 0.005% or less), Ti forms an oxide, acts as an intragranular ferrite formation nucleus in the HAZ, and has an effect of refining the HAZ structure. In order to exhibit such an effect of TiN, it is necessary to add at least 0.005% Ti. However, if the amount of Ti is too large, TiN coarsening and precipitation hardening due to TiC occur and the low temperature toughness is deteriorated, so the upper limit was limited to 0.030%.
[0017]
Al is an element usually contained in steel as a deoxidizing material, and has an effect on refinement of the structure. However, if the amount of Al exceeds 0.06%, Al-based non-metallic inclusions increase to impair the cleanliness of the steel, so the upper limit was made 0.06%. However, deoxidation can be performed with Ti or Si, and Al need not necessarily be added.
The above are the main components of the steel pipe base material of the present invention, but the following components are selectively contained as necessary.
[0018]
B is a very effective element in order to dramatically improve the hardenability of steel in a very small amount and to obtain the target martensite-based structure. Further, B enhances the hardenability improvement effect of Mo, and synergistically increases the hardenability by coexisting with Nb. On the other hand, if added excessively, not only the low-temperature toughness is deteriorated, but also the effect of improving the hardenability of B may be lost, so the upper limit was made 0.0030%.
[0019]
N forms TiN and suppresses coarsening of the austenite grains of HAZ during reheating of the slab and improves the low temperature toughness of the base material and HAZ. The minimum amount required for this is 0.001%. However, if the amount of N is too large, it will cause deterioration of the HAZ toughness due to slab surface defects and solute N, and decrease in the effect of improving the hardenability of B, so the upper limit must be limited to 0.006%.
[0020]
Next, the purpose of adding V, Cu, Cr, Ca, REM, and Mg will be described.
The main purpose of adding these elements to the basic components of the steel pipe base material of the present invention is to further improve the strength and toughness and expand the steel material size that can be manufactured without impairing the excellent characteristics of the steel of the present invention. It is to measure. Therefore, the amount of addition is naturally limited.
[0021]
V has almost the same effect as Nb, but the effect is weaker than Nb. However, the effect of V addition in the ultra high strength steel is large, and the combined addition of Nb and V makes the excellent characteristics of the steel of the present invention more remarkable. The upper limit is acceptable up to 0.10% from the viewpoint of HAZ toughness and field weldability, but the addition of 0.03 to 0.08% is particularly desirable.
[0022]
Cu increases the strength of the base metal and the welded portion, but if too much, the HAZ toughness and on-site weldability are remarkably deteriorated. For this reason, the upper limit of the amount of Cu is 1.0%.
Cr increases the strength of the base metal and the welded portion, but if too much, the HAZ toughness and on-site weldability are significantly deteriorated. For this reason, the upper limit of Cr amount is 0.6%.
Ca and REM control the form of sulfide (MnS) and improve low-temperature toughness (such as an increase in absorbed energy in the Charpy test). When Ca content is 0.006% and REM exceeds 0.02%, a large amount of CaO-CaS or REM-CaS is formed, resulting in large clusters and large inclusions, not only harming the cleanliness of the steel, It also adversely affects on-site weldability. For this reason, the upper limit of the Ca addition amount is limited to 0.006% or the REM addition amount condition is limited to 0.02%. In ultra-high strength line pipes, the S and O contents are reduced to 0.001% and 0.002% or less, respectively, and ESSP (Effestive Sulphide Shape Controlling) is an index for controlling the shape of MnS clusters shown below. It is particularly effective to adjust Ca, S, and O so that (Parameter) satisfies 0.5 ≦ ESSP ≦ 10.0.
[0023]
ESSP = (Ca) [1-124 (O)] / 1.25S (1)
When the above ESSP is less than 0.5, a large amount of CaO-CaS is formed, resulting in coarse clusters and coarse inclusions, which deteriorate weldability such as weld cracks. When the ESSP exceeds 10.0, the shape of MnS Since the effect of control is lost, ESSP is specified to be 0.5 to 10.0.
[0024]
Mg forms finely dispersed oxides and suppresses the coarsening of the weld heat affected zone to improve the low temperature toughness. If it is 0.006% or more, a coarse oxide is produced and the toughness is deteriorated.
In addition to the limitation of the individual additive elements described above, it is desirable to further limit P, which is an index representing the hardenability described below, to 1.9 ≦ P ≦ 4.0.
[0025]
P = 2.7C + 0.4Si + Mn + 0.8Cr + 0.45 (Ni + Cu) + (1 + β) Mo-1 + β (2)
However, β = 1 when B ≧ 3 ppm, and β = 0 when B <3 ppm.
The reason for controlling P as described above is to achieve the desired balance between strength and low temperature toughness. The lower limit of the P value is set to 1.9 in order to obtain a strength of 900 MPa or more and excellent low temperature toughness. The upper limit of the P value is set to 4.0 in order to maintain excellent HAZ toughness and on-site weldability.
[0026]
The above is the grounds for limiting the components contained in the steel pipe base material of the present invention. Even if the above chemical components are present, a fine martensite + bainite-based structure, which is the structure of the present invention, can be obtained. Therefore, the desired characteristics cannot be obtained unless the manufacturing conditions are appropriate. The principle method of obtaining a fine martensite-based structure is to process the recrystallized grains in the non-recrystallized temperature range to obtain austenite grains flattened in the thickness direction, which exceeds the critical cooling rate at which ferrite formation is suppressed. It is to cool at the cooling rate.
[0027]
A desirable manufacturing method is to reheat a steel slab having the chemical component of the present invention to 950 to 1250 ° C. and roll at a steel material temperature of 700 ° C. or higher so that the cumulative reduction amount at 700 to 950 ° C. is 50% or higher. Then, it cools to 550 degrees C or less with the cooling rate of 10 degrees C or more. A if necessary C1 Tempering is performed at a temperature below the transformation point.
In the steel pipe of the present invention, the steel plate thus manufactured is formed into a tubular shape, and then the butted portions at both ends of the steel plate are arc-welded and further expanded to form a steel pipe.
[0028]
Next, the reasons for limiting the components of the weld metal part of the steel pipe of the present invention will be described.
The amount of C is limited to 0.03 to 0.14%. Carbon is extremely effective in improving the strength of steel, and at least 0.03% is necessary to obtain the target strength in the martensite structure. However, if the amount of C is too large, cold cracking is likely to occur, and the maximum hardness of the HAZ of the so-called T-cross portion where the local weld and seam welding intersect increases, so the upper limit was made 0.14%. Furthermore, the upper limit is preferably 0.10%.
[0029]
Si needs to be 0.05% or more to prevent blowholes, but if the content is large, the low temperature toughness is remarkably deteriorated, so the upper limit was made 0.40%. In particular, when performing inner and outer surface welding or multilayer welding, the low temperature toughness of the reheated portion is deteriorated.
Mn is an indispensable element for securing an excellent balance between strength and low temperature toughness, and its lower limit is 1.2%. However, too much Mn not only promotes segregation and deteriorates low-temperature toughness, but also makes it difficult to produce a welding material, so the upper limit was made 2.2%.
[0030]
For P and S, it is desirable that the amount of P and S is low in order to reduce the low temperature toughness and the low temperature cracking susceptibility, and the upper limit is defined as 0.010%.
The purpose of adding Ni is to increase the hardenability to ensure the strength and to further improve the low temperature toughness. If it is 1.3% or less, it is difficult to obtain the target strength and low temperature toughness. On the other hand, if the content is too high, there is a risk of hot cracking, so the upper limit was made 3.2%.
[0031]
Although the difference in the effects of Cr, Mo, and V cannot be strictly discriminated, any of them is added to obtain high strength by enhancing the hardenability. If the total amount of one or more of Cr, Mo and V is 1.0% or less, the effect is not sufficient. On the other hand, if added in a large amount, the risk of cold cracking increases, so the upper limit was made 2.5%.
B is an element that enhances the hardenability in a small amount and is effective for improving the low temperature toughness of the weld metal. However, if the content is too large, the low temperature toughness is lowered, so the content range is set to 0.005% or less.
[0032]
The weld metal may contain, as other components, elements such as Ti, Al, Zr, Nb, and Mg that are added as necessary in order to improve refining and solidification during welding. The balance is iron and inevitable impurities.
The ultra-high-strength steel sheet of the present invention is manufactured by casting the steel having the above-mentioned components, hot working, and then quenching or tempering in some cases. In order to achieve ultra-high strength with a tensile strength of 900 MPa or more, it is necessary to suppress the formation of ferrite by making steel a microstructure mainly composed of low-temperature transformation structure such as martensite and bainite.
[0033]
The weld metal is a solidified structure after welding, and in the weld metal having a slow cooling rate, the present invention obtains the above-mentioned target strength, and further obtains excellent low temperature toughness similar to the steel sheet of the present invention. It is necessary to adjust the chemical composition and structure of the metal.
Ni enhances hardenability and makes it possible to obtain high strength even at a low cooling rate, and promotes formation of retained austenite between martensite laths and improves low temperature toughness.
[0034]
In the present invention, desired strength and low-temperature toughness can be obtained by making the Ni content of the weld metal 1% or more higher than the steel plate component and making the weld metal part and the weld heat affected zone a bainite martensite structure. When the amount of Ni in the weld metal is 1% lower than the steel plate component, the above effect cannot be obtained. Therefore, in the present invention, the lower limit is set to 1%.
Next, a method for producing the steel pipe of the present invention will be described.
[0035]
The line pipe aimed by the present invention is usually a size with a diameter of 450 mm to 1500 mm and a wall thickness of about 10 mm to 40 mm. As a method of manufacturing such a steel pipe with a high rate, pipes are manufactured in a UO process in which a steel sheet is formed into a U shape and then an O shape. After the tack welding is performed, a steel pipe manufacturing method is established in which the tack welded portion is subjected to main welding from the inner and outer surfaces by submerged arc welding or the like and then expanded to increase the roundness.
[0036]
An arc welding method such as submerged arc welding is welding with a large dilution of the base material, and in order to obtain desired characteristics, that is, a weld metal composition, it is necessary to select a welding material in consideration of the dilution of the base material.
Below, the reason for limitation of the chemical composition of the welding wire used for welding at the time of manufacturing the ultra high strength line pipe of the present invention will be described. In addition, since the tack welding of the steel plate mating part performed before the main welding of the present invention has a small welding area and the influence of the quality of the weld metal part is small compared to the main welding, the components of the welding wire used for the main welding are All are defined as follows, but the components of the welding wire used for tack welding need not be particularly specified for components other than C, Si, and Mn.
[0037]
C was set to 0.01 to 0.12% in consideration of dilution with the base material component and mixing of C from the atmosphere in order to obtain a range of C amount required for the weld metal.
In order to obtain a range of Si amount required for the weld metal, Si is set to 0.3% or less in consideration of dilution by the base material component.
Mn is set to 1.2% to 2.4% in consideration of dilution by a base material component in order to obtain a range of Mn amount required for the weld metal.
[0038]
In order to obtain the range of Ni amount required for the weld metal, Ni is set to 4.0% to 8.5% in consideration of dilution by the base material component.
In order to obtain a total amount range of one or more of Cr, Mo, V required for the weld metal, Cr, Mo, V is 3.0 in consideration of dilution by the base material component. % To 5.0%.
[0039]
In addition, it is desirable that impurities of P and S are as small as possible, and B can be added to ensure strength. Ti, Al, Zr, Nb, Mg, etc. are used for the purpose of deoxidation.
The tack welding and the main welding of the present invention can be performed not only with a single electrode but also with a plurality of electrodes. In the case of welding with a plurality of electrodes, it is possible to combine various wires, and it is not necessary for each wire to be in the above-mentioned component range, and the average composition from each wire component and consumption may be in the above-described component range.
[0040]
Flux used for main welding such as submerged arc welding can be roughly classified into firing-type flux and fusion-type flux. Firing-type fluxes have the advantage that alloy materials can be added and the amount of diffusible hydrogen is low, but they have the disadvantage of being easily pulverized and difficult to use repeatedly. On the other hand, the melt-type flux is in the form of glass powder, has the advantages of high grain strength, hardly absorbs moisture, and has the disadvantage that diffusible hydrogen is slightly high. In the case of ultra-high strength such as the present invention, cold cracking is likely to occur. From this point, a firing mold is desirable. On the other hand, a molten mold that can be recovered and used repeatedly has the advantage of low cost for mass production. is there. The cost is high in the baking mold, and the necessity of strict quality control is a problem in the melting mold, but it is within a range that can be handled industrially, and either can be used essentially.
[0041]
The preferred ranges for welding conditions are as follows.
The tack welding performed first may be MAG arc welding, MIG arc welding, or TIG arc welding. Usually, MAG arc welding. Next, the main welding performed after tack welding is usually submerged arc welding, but may be TIG arc welding, MIG arc welding, or MAG arc welding. An appropriate range of the welding speed is about 1 to 3 m / min. Welding less than 1 m / min is inefficient as line pipe seam welding, and the bead shape is not stable at high speed welding exceeding 3 m / min. If tack welding and subsequent main welding overlap, it is preferable that welding heat input is as low as possible. Further, the arc welding of the main welding may be performed with any number of passes. Although the welding heat input varies depending on the plate thickness, for example, in the case of a plate thickness of 16 mm, it is desirable that the welding heat input be 1.0 to 2.7 kJ / mm. If the heat input is too small, the penetration becomes insufficient, the number of weldings increases, the work efficiency deteriorates, and if the welding heat input is too large, the heat-affected zone is greatly softened and the toughness of the welded portion is also reduced.
[0042]
After these seam welding, roundness is improved by pipe expansion. In order to make a perfect circle, it is necessary to deform to the plastic region. However, in the case of high strength steel such as the present invention, a tube expansion ratio of about 0.7% or more (= (circumference after tube expansion−circumference before tube expansion) / (Circumference before pipe expansion) is necessary, but if large pipe expansion exceeding 2% is performed, the toughness deterioration due to plastic deformation increases in both the base metal and the welded part, so the pipe expansion ratio should be 0.7-2% or less. Is desirable.
[0043]
【Example】
Examples of the present invention and effects thereof will be specifically described below.
Invented steels (A steel to D steel) with components satisfying the scope of the present invention shown in Table 1 and comparative steels (E steel, F steel) with components outside the scope of the present invention are melted in a 300-ton converter, then continuously cast steel After being reheated to 1100 ° C. and rolled in the recrystallization region, controlled rolling is performed until the cumulative reduction amount of 900 to 750 ° C. becomes 75% up to 16 mm, and then the water cooling stop temperature becomes 200 to 450 ° C. The steel sheet was manufactured by water cooling as described above. As a result, as shown in Table 1, the strength of the steel sheets of the invention steels (A steel to D steel) is the target range (900 MPa or more) of the present invention, and low temperature toughness (energy absorbed at −30 ° C. in Charpy test): 230J or higher) was also high. On the other hand, the strength of the steel plate of E steel with a high C content and no added Ni is within the target range of the present invention, but the low temperature toughness is low, and the low temperature toughness of the steel plate of F steel with a low C content is within the target range. However, the strength is low.
[0044]
[Table 1]
[Table 2]
These steel sheets are further formed into a tubular shape at the UO factory, and the mating part of the steel sheets is 80% Ar + 20% CO. 2 After performing tack welding using MAG arc welding with a shielding gas of 3mm, using the welding wire and flux shown in Table 2, provisional welding under the welding conditions of 3 electrodes, 2.0m / min, heat input 1.5KJ / mm The inner and outer surfaces of the part were subjected to main welding by submerged arc welding of one pass each, and thereafter, the pipe was expanded at a pipe expansion rate of 1%. Table 3 shows the results of evaluating the properties of the obtained steel pipe.
[0047]
In the invention examples (Examples Nos. 1 to 6) welded using steel and welding wires having components satisfying the scope of the present invention shown in Table 2, a good weld bead is obtained in the steel pipe seam weld, and the chemistry of the weld metal part The components satisfy the scope of the present invention, and the weld metal strength (900 MPa or more), the difference between the tensile strength of the weld metal and the tensile strength of the steel sheet, and the appropriate range (-100 MPa or more) are appropriate. The distance (Δd) between the metal part and the outer surface weld metal part (Δd: more than 0 mm)) was also appropriate. In addition, the steel pipes of the inventive examples satisfying these scopes of the present invention have mechanical properties such as the target strength and low temperature toughness of the base metal part and the welded part, and the pipe fracture can be achieved even in the burst test. .
[0048]
On the other hand, the implementation No. of the comparative example. 7 to 9, although the base material component is within the scope of the present invention, the wire component is outside the scope of the present invention (No. 7: Ni is low, No. 8: C is high, No. 9: Ni is Therefore, the components of the weld metal part deviated from the scope of the present invention. As a result, no. In No. 7, the strength of the weld metal was lowered, and the difference between the tensile strength of the weld metal and the tensile strength of the steel plate was out of the proper range (−100 MPa or more), so that the weld fracture occurred in the burst test. No. In No. 8, a cold crack occurred in the welded part. Since hot cracking occurred in No. 9, the tensile test and the burst test could not be performed.
[0049]
Comparative Example No. No. 10, the component of the welding wire is within the range of the present invention, but the component of the steel sheet is outside the range of the present invention, so the low temperature toughness of the steel pipe base material did not reach the target (low level).
However, no. 11 and no. 12, the components of the base metal and the welding wire are within the scope of the present invention, but the interval (Δd) between the inner surface weld metal part and the outer surface weld metal part in the weld part formed by the main welding is within the scope of the present invention. Since (Δd> 0) was deviated, the weld heat affected zone fracture occurred in the burst test. Moreover, No. of the comparative example. No. 13 does not satisfy the required characteristics of the line pipe because the weld metal and weld metal C is out of the scope of the present invention (high), and the weld metal has low toughness. In Comparative Example No14, the base metal and welding wire components are within the scope of the present invention, but the difference between the tensile strength of the weld metal and the tensile strength of the steel plate deviated from the appropriate range (-100 MPa or more), so in the burst test, the welding heat Affected zone fracture occurred.
[0050]
[Table 3]
[Table 4]
[Table 5]
【The invention's effect】
According to the present invention, it is possible to realize an ultra-high strength line pipe having a tensile strength of 900 MPa or more (API standard X100 or more) that has excellent balance of low-temperature toughness in both the base material and the welded portion and that can be easily welded on site. Pipeline laying costs are reduced, which can contribute to solving global energy problems.
[Brief description of the drawings]
FIG. 1 is a cross-sectional view of a steel pipe seam weld according to the present invention.
FIG. 2 is a cross-sectional view of a conventional steel pipe seam weld.
[Explanation of symbols]
1 ... Outer weld metal part of this weld metal part
2. Internal weld metal part of the main weld metal part
3. Base material part of steel pipe
4 ... Tack weld metal part
Claims (5)
Ｃ：０．０３〜０．１０％、
Ｓｉ：０．６％以下、
Ｍｎ：１．７〜２．５％、
Ｐ：０．０１５％以下、
Ｓ：０．００３％以下、
Ｎｉ：０．１〜１．０％、
Ｍｏ：０．１５〜０．６０％、
Ｎｂ：０．０１〜０．１０％、
Ｔｉ：０．００５〜０．０３０％、
Ａｌ：０．０６％以下、
を含有し、さらに重量％で、Ｂ：０．００３０％以下、Ｎ：０．００１〜０．００６％、Ｖ：０．１０％以下、Ｃｕ：１．０％以下、Ｃｒ：１．０％以下、Ｃａ：０．０１％以下、ＲＥＭ：０．０２％以下、Ｍｇ：０．００６％以下の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる母材部と、
重量％で、
Ｃ：０．０３〜０．１４％、
Ｓｉ：０．０５〜０．４０％、
Ｍｎ：１．２〜２．２％、
Ｐ：０．０１０％以下、
Ｓ：０．０１０％以下、
Ｎｉ：１．３〜３．２％、
Ｃｒ、Ｍｏ、Ｖのうちの１種または２種以上の合計量が１．０〜２．５％、
Ｂ：０．００５％以下、
を含有し、残部が鉄および不可避的不純物からなる溶接金属部からなり、かつ、溶接金属部のＮｉ量が母材部のＮｉ量に比べて１％以上高く、溶接金属部及び母材の溶接熱影響部を含むシーム溶接部の組織がベイナイト・マルテンサイトからなることを特徴とする請求項１に記載のシーム溶接部の低温靱性の優れた超高強度溶接鋼管。% By weight
C: 0.03-0.10%,
Si: 0.6% or less,
Mn: 1.7-2.5%,
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 1.0%,
Mo: 0.15-0.60%,
Nb: 0.01-0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less,
In addition, by weight%, B: 0.0030% or less, N: 0.001 to 0.006%, V: 0.10% or less, Cu: 1.0% or less, Cr: 1.0% Hereinafter, Ca: 0.01% or less, REM: 0.02% or less, Mg: 0.006% or less of one or two or more of the base material part consisting of iron and inevitable impurities,
% By weight
C: 0.03-0.14%,
Si: 0.05-0.40%,
Mn: 1.2-2.2%,
P: 0.010% or less,
S: 0.010% or less,
Ni: 1.3-3.2%
The total amount of one or more of Cr, Mo and V is 1.0 to 2.5%,
B: 0.005% or less,
And the balance is made of a weld metal part made of iron and inevitable impurities, and the amount of Ni in the weld metal part is 1% or more higher than the amount of Ni in the base metal part. The super high strength welded steel pipe with excellent low temperature toughness of the seam weld according to claim 1, wherein the structure of the seam weld including the heat affected zone is composed of bainite martensite.
Ｃ：０．０３〜０．１０％、
Ｓｉ：０．６％以下、
Ｍｎ：１．７〜２．５％、
Ｐ：０．０１５％以下、
Ｓ：０．００３％以下、
Ｎｉ：０．１〜１．０％、
Ｍｏ：０．１５〜０．６０％、
Ｎｂ：０．０１〜０．１０％、
Ｔｉ：０．００５〜０．０３０％、
Ａｌ：０．０６％以下、
を含有し、さらに重量％で、Ｂ：０．００３０％以下、Ｎ：０．００１〜０．００６％、Ｖ：０．１０％以下、Ｃｕ：１．０％以下、Ｃｒ：１．０％以下、Ｃａ：０．０１％以下、ＲＥＭ：０．０２％以下、Ｍｇ：０．００６％以下の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼板の両端部を付き合わせた後、該付き合わせ部を、重量％で、Ｃ：０．０１〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１．２〜２．４% を含有しＦｅを主成分とする溶接ワイヤーを用いて、外面から仮付け溶接を行った後、該仮付け溶接部を、重量％で、Ｃ：０．０１〜０．１２％、Ｓｉ：０．３％以下、Ｍｎ：１．２〜２．４% 、Ｎｉ：４．０〜８．５％、Ｃｒ、Ｍｏ、Ｖの１種又は２種以上の合計量３．０〜５．０％を含有し、かつＮｉ量が前記鋼板のＮｉ量に比べて１％以上高いＦｅを主成分とする溶接ワイヤーおよびフラックスを用いて、溶接によって形成される内面溶接金属部と外面溶接金属部の間隔が０ｍｍ超であり、かつ、内面溶接金属部と外面溶接金属部が前記仮付け溶接によって形成される仮付け溶接金属部とそれぞれ重複するように、前記仮付け溶接部を内面及び外面から本溶接を行うことを特徴とするシーム溶接部の低温靱性の優れた超高強度溶接鋼管の製造方法。% By weight
C: 0.03-0.10%,
Si: 0.6% or less,
Mn: 1.7-2.5%,
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 1.0%,
Mo: 0.15-0.60%,
Nb: 0.01-0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less,
In addition, by weight%, B: 0.0030% or less, N: 0.001 to 0.006%, V: 0.10% or less, Cu: 1.0% or less, Cr: 1.0% Hereinafter, both ends of a steel sheet containing one or more of Ca: 0.01% or less, REM: 0.02% or less, Mg: 0.006% or less, the balance being iron and unavoidable impurities. After the attachment, the attachment part contains C: 0.01 to 0.12%, Si: 0.3% or less, and Mn: 1.2 to 2.4% by weight%. After performing tack welding from the outer surface using a welding wire as a component, the tack welded portion is expressed in terms of% by weight, C: 0.01 to 0.12%, Si: 0.3% or less, Mn : 1.2 to 2.4%, Ni: 4.0 to 8.5%, containing one or more of Cr, Mo, V or a total amount of 3.0 to 5.0%, or The distance between the inner weld metal part and the outer weld metal part formed by welding is greater than 0 mm using a welding wire and a flux mainly composed of Fe whose Ni content is 1% or more higher than the Ni content of the steel sheet. And performing the main welding from the inner surface and the outer surface so that the inner weld metal portion and the outer weld metal portion overlap with the tack weld metal portion formed by the tack welding, respectively. A method for producing an ultra-high-strength welded steel pipe excellent in low temperature toughness of a seam weld.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP29422599A JP3814112B2 (en) | 1999-10-15 | 1999-10-15 | Super high strength steel pipe excellent in low temperature toughness of seam welded portion and manufacturing method thereof |
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| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP29422599A JP3814112B2 (en) | 1999-10-15 | 1999-10-15 | Super high strength steel pipe excellent in low temperature toughness of seam welded portion and manufacturing method thereof |
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| Publication Number | Publication Date |
|---|---|
| JP2001113374A JP2001113374A (en) | 2001-04-24 |
| JP3814112B2 true JP3814112B2 (en) | 2006-08-23 |
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| JP29422599A Expired - Fee Related JP3814112B2 (en) | 1999-10-15 | 1999-10-15 | Super high strength steel pipe excellent in low temperature toughness of seam welded portion and manufacturing method thereof |
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Cited By (2)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN102676744A (en) * | 2012-06-18 | 2012-09-19 | 北京科技大学 | Technique for refined production of antiacid pipeline steel through VD-LF-VD |
| CN102725098A (en) * | 2010-01-27 | 2012-10-10 | 株式会社神户制钢所 | Welding metal having excellent low-temperature toughness and drop-weight characteristics |
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| WO2005108636A1 (en) * | 2004-05-11 | 2005-11-17 | Sumitomo Metal Industries, Ltd. | Super high strength uoe steel pipe and method for production thereof |
| JP4655670B2 (en) * | 2005-02-24 | 2011-03-23 | Ｊｆｅスチール株式会社 | Manufacturing method of high strength welded steel pipe with low yield ratio and excellent weld toughness |
| JP5055899B2 (en) * | 2006-08-30 | 2012-10-24 | Ｊｆｅスチール株式会社 | Method for producing high-strength welded steel pipe excellent in weld heat-affected zone toughness and having tensile strength of 760 MPa or more, and high-strength welded steel pipe |
| JP5251089B2 (en) * | 2006-12-04 | 2013-07-31 | 新日鐵住金株式会社 | Welded steel pipe for high-strength thick-walled line pipe excellent in low-temperature toughness and manufacturing method |
| JP2013078775A (en) | 2011-10-03 | 2013-05-02 | Jfe Steel Corp | Welded steel pipe excelling in toughness of welding heat affected part, and method for manufacturing the same |
| CN103131965B (en) * | 2011-11-25 | 2015-10-28 | 中国石油天然气集团公司 | A kind of for based on the X100 gross distortion pipe line steel of stress design and preparation method |
| KR101428201B1 (en) | 2012-10-30 | 2014-08-07 | 주식회사 포스코 | Submerged arc welded joint |
| KR101449209B1 (en) * | 2012-12-27 | 2014-10-08 | 주식회사 포스코 | Welding joint |
| JP6447253B2 (en) * | 2015-03-06 | 2019-01-09 | 新日鐵住金株式会社 | High strength steel for welding |
-
1999
- 1999-10-15 JP JP29422599A patent/JP3814112B2/en not_active Expired - Fee Related
Cited By (2)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN102725098A (en) * | 2010-01-27 | 2012-10-10 | 株式会社神户制钢所 | Welding metal having excellent low-temperature toughness and drop-weight characteristics |
| CN102676744A (en) * | 2012-06-18 | 2012-09-19 | 北京科技大学 | Technique for refined production of antiacid pipeline steel through VD-LF-VD |
Also Published As
| Publication number | Publication date |
|---|---|
| JP2001113374A (en) | 2001-04-24 |
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