JP5348383B2 - High toughness welded steel pipe with excellent crushing strength and manufacturing method thereof - Google Patents
High toughness welded steel pipe with excellent crushing strength and manufacturing method thereof Download PDFInfo
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Abstract
Description
本発明は、石油や天然ガスの輸送に使用される高強度高靱性ラインパイプ用鋼管およびその製造方法に関し、特に厚鋼板を冷間で成型し溶接して製造される圧潰強度に優れた高靱性溶接鋼管およびその製造方法。 The present invention relates to a steel pipe for a high strength and high toughness line pipe used for transportation of oil and natural gas and a method for producing the same, and in particular, a high toughness with excellent crushing strength produced by cold forming and welding a thick steel plate. Welded steel pipe and manufacturing method thereof.
海底パイプラインなどの高い外圧環境化で使用されるラインパイプは、外圧による高い圧縮応力が負荷されることにより圧潰する危険性がある。そのため、海底パイプラインで使用されるラインパイプは、造管した状態で十分な圧縮強度を有する必要がある。しかしながら、ＵＯＥ鋼管のように厚鋼板を冷間加工した後、拡管することで造管される鋼管の場合、最終工程である拡管で大きな引張負荷を受ける。したがって、それらの鋼管の圧縮降伏応力は、引張負荷時に発生した背応力により鋼管の引張降伏応力よりも低下することになる。 Line pipes used in high external pressure environments such as submarine pipelines have a risk of being crushed when subjected to high compressive stress due to external pressure. Therefore, the line pipe used in the submarine pipeline needs to have a sufficient compressive strength in a piped state. However, in the case of a steel pipe that is made by cold-working a thick steel plate like a UOE steel pipe and then expanding it, the pipe is subjected to a large tensile load in the final process. Therefore, the compressive yield stress of these steel pipes is lower than the tensile yield stress of the steel pipes due to the back stress generated during the tensile load.
よって、鋼管の圧縮強度を確保するためには、厚鋼板の設計強度を高く設定する必要がある。強度を強化元素で補う場合には、合金コストの増大や母材および溶接熱影響部の靱性劣化が懸念されるため、造管後の圧縮降伏応力の低下が少ない厚鋼板もしくは造管方法の開発が求められている。また、海底に敷設するラインパイプ母材部には、低温での優れた脆性き裂伝播停止性能が要求される。このため、圧潰強度と脆性き裂伝播停止性能を両立するラインパイプの開発が求められている。 Therefore, in order to ensure the compressive strength of the steel pipe, it is necessary to set the design strength of the thick steel plate high. When strength is supplemented with strengthening elements, there is concern about increased alloy costs and toughness deterioration of the base metal and weld heat-affected zone, so the development of a thick steel plate or pipe making method with less reduction in compressive yield stress after pipe making Is required. In addition, the line pipe preform laid on the seabed is required to have excellent brittle crack propagation stopping performance at low temperatures. For this reason, development of a line pipe that achieves both crushing strength and brittle crack propagation stopping performance is required.
このような要求に対し、特許文献１および特許文献２では、造管時のＯプレス圧縮率と拡管率をパラメータに、圧縮率／拡管率を最適な範囲まで低減することによって、造管後における鋼管の圧縮降伏応力の低下を抑制する方法が開示されている。たとえば、特許文献２には、Ｏ成形時のアプセット率（すなわち圧縮率）αと拡管時の拡管率βとの比をα／β≧０．３５とする技術が開示されている。また、特許文献２では、拡管率を極めて大きくすることにより、造管後における鋼管の圧縮降伏応力の低下を抑制する方法も開示されている。特許文献３では、縮管と拡管の順序と程度を最適化することによって、外圧による鋼管の圧潰強度を向上させる方法が開示されている。特許文献４から７には、造管後に熱処理、もしくはコーティング加熱による低温ひずみ時効により、造管工程で鋼管に付与された背応力を低減することにより、鋼管の圧縮降伏応力の低下を抑制する方法が開示されている。
しかし、特許文献１および特許文献２で示されているような最適な圧縮率／拡管率に造管条件を設定するためには、Ｏプレス圧縮率を通常よりも極めて大きくする必要がある。Ｏプレスの圧縮率を増大させることは、Ｏプレス機のプレス能力を増強する必要があり、新規設備導入や設備改修によるコストの増大が問題となる。 However, in order to set the pipe making conditions to the optimum compression ratio / pipe expansion ratio as shown in Patent Document 1 and Patent Document 2, it is necessary to make the O-press compression ratio much larger than usual. Increasing the compression rate of the O press requires an increase in the press capability of the O press machine, and increases the cost due to the introduction of new equipment and the repair of equipment.
さらに、圧縮強度の確保が問題となる海底パイプライン用ラインパイプは、耐座屈性能確保の観点から厚肉で設計されることが多く、このことはＯプレスの圧縮率を増大させることとなる。また、拡管率を低下させることにより、最適な範囲にすることもできるが、鋼管の真円度を低下させることとなり、このことも耐座屈性能の観点から厳しい真円度の要求がなされる海底パイプライン用ラインパイプを製造する上においては限界がある。 Furthermore, line pipes for submarine pipelines, where securing compressive strength is a problem, are often designed with a thick wall from the viewpoint of securing buckling resistance, which increases the compressibility of the O-press. . In addition, the pipe expansion ratio can be reduced to the optimum range, but the roundness of the steel pipe will be lowered, and this also requires severe roundness from the viewpoint of buckling resistance. There are limitations in manufacturing line pipes for submarine pipelines.
また、特許文献２および３に記載のように、拡管率を極めて大きくすることや縮管と拡管とを行うことは、過度な加工硬化による表面硬さの上昇や、残留応力の増大による脆性き裂伝播停止性能の劣化が懸念される。 Further, as described in Patent Documents 2 and 3, increasing the tube expansion rate or performing tube contraction and tube expansion is not brittle due to an increase in surface hardness due to excessive work hardening and an increase in residual stress. There is concern about degradation of crack propagation stopping performance.
また、特許文献４から７に記載のように、造管後のコーティング加熱条件を最適化することにより，低温ひずみ時効処理を行うことは、圧縮降伏応力の低下を抑制するという観点では絶大な効果があるが、鋼管の引張の応力-ひずみ曲線が造管後のラウンドハウス型からリューダース型に変わり、曲げ座屈性能などの鋼管の変形能を低下させる。 Moreover, as described in Patent Documents 4 to 7, performing the low-temperature strain aging treatment by optimizing the coating heating conditions after pipe forming is a great effect from the viewpoint of suppressing the decrease in compressive yield stress. However, the tensile stress-strain curve of the steel pipe changes from the round house type after pipe making to the Luders type, which lowers the deformability of the steel pipe such as bending buckling performance.
さらに、コーティング加熱の条件は、使用するコーティング材によって変わり、必ずしも狙いとするコーティング加熱条件に合致させることができるとは限らず、コーティング加熱のかわりに熱処理によって低温ひずみ時効処理を行う場合は、工程が増えることにより生産性を著しく損なうこととなる。 Furthermore, the coating heating conditions vary depending on the coating material used, and may not necessarily match the target coating heating conditions. When performing low-temperature strain aging treatment by heat treatment instead of coating heating, As a result, the productivity will be significantly impaired.
上述したように、従来の技術では溶接性の低下、変形性能の低下、生産性の低下や脆性き裂伝播停止性能の低下を生じることなく、圧潰強度に優れた溶接鋼管を製造することは、困難であった。 As described above, in the conventional technology, producing a welded steel pipe excellent in crushing strength without causing degradation in weldability, degradation in deformation performance, degradation in productivity and degradation in brittle crack propagation stopping performance, It was difficult.
そこで、本研究では、真円度および変形性能を低下させることなく、高生産性、低コストで製造でき、優れた脆性き裂伝播停止性能を有する圧潰強度に優れた溶接鋼管およびその製造方法を提供することを目的とする。 Therefore, in this study, a welded steel pipe that can be manufactured at high productivity and low cost without reducing roundness and deformation performance, has excellent brittle crack propagation stopping performance, and a manufacturing method thereof. The purpose is to provide.
発明者らは、前記の課題を解決するために、鋼板のミクロ組織およびミクロ組織を達成するための製造方法、特に制御圧延、加速冷却とその後の再加熱という製造プロセスについて鋭意検討し、以下の知見を得た。 In order to solve the above-mentioned problems, the inventors diligently studied the microstructure of the steel sheet and the manufacturing method for achieving the microstructure, particularly the manufacturing process of controlled rolling, accelerated cooling and subsequent reheating, and the following. Obtained knowledge.
まず、優れた脆性き裂伝播停止性能を得るためには、二相域での圧延によりフェライトを加工し、圧延面での（１００）面の集合組織を発達させ、脆性き裂伝播時にセパレーションを発生させることが必要不可欠であることがわかった。また、管厚中心位置での圧延面の（１００）面の集積度を１．５以上とし、なおかつ主な金属組織であるフェライト相とベイナイト相との硬度差を５０以上とすることで、脆性破壊伝播停止性能の評価法であるＤＷＴＴ試験（Drop Weight Tear Test；落重試験）を行った際に、破面にセパレーションが発生し、より低温まで高い延性破面率を確保できることがわかった。 First, in order to obtain excellent brittle crack propagation stopping performance, ferrite is processed by rolling in a two-phase region, a (100) plane texture is developed on the rolled surface, and separation is performed during brittle crack propagation. It turns out that it is essential to generate. Further, the degree of integration of the (100) plane of the rolled surface at the tube thickness center position is set to 1.5 or more, and the hardness difference between the ferrite phase and the bainite phase, which is the main metal structure, is set to 50 or more. It was found that when the DWTT test (Drop Weight Tear Test), which is an evaluation method of the fracture propagation stopping performance, was performed, separation occurred on the fracture surface, and a high ductile fracture surface ratio could be secured even at a lower temperature.
次に、発明者らは、前記課題を解決するために、鋼管の圧潰強度の指標となるパラメータである圧縮降伏応力の高い鋼管組織形態ならびに圧縮降伏応力を向上させる鋼管素材および鋼管の製造方法について検討した。 Next, in order to solve the above problems, the inventors of the present invention relate to a steel pipe structure that has a high compressive yield stress and a steel pipe material that improves the compressive yield stress, and a method for manufacturing the steel pipe, which are parameters that serve as indices for the crushing strength of the steel pipe. investigated.
まず、鋼管素材の製造方法として一般的である制御圧延、加速冷却という製造プロセスにより種々の組織形態に作りこんだ鋼管素材を、同一条件で造管して溶接鋼管を製造した。 First, welded steel pipes were produced by producing steel pipe materials that were formed into various structural forms by the manufacturing processes of controlled rolling and accelerated cooling, which are common methods for producing steel pipe materials, under the same conditions.
次に溶接鋼管からＬ方向の全厚引張試験片およびＬ方向の圧縮試験片を採取し、引張降伏応力に対する圧縮降伏応力の比（圧縮降伏応力／引張降伏応力）を比較した結果、フェライト＋ベイナイト組織鋼に代表される複相組織鋼管よりもベイナイト単相組織鋼管の方が、この比が大きいことがわかった。 Next, full thickness tensile specimens in the L direction and compression specimens in the L direction were taken from the welded steel pipe, and the ratio of compressive yield stress to tensile yield stress (compressive yield stress / tensile yield stress) was compared. It was found that this ratio was larger in the bainite single-phase steel pipe than in the double-phase steel pipe represented by the microstructure steel.
さらに、ベイナイト単相組織鋼管のミクロ組織中の硬質第２相である島状マルテンサイト（以下Ｍ−Ａという）を減少させることで，造管段階での硬質相周辺で発生する局所的なひずみ勾配を緩和し、バウシンガー効果による圧縮降伏応力の低下を抑制することが可能であることがわかった。 Furthermore, by reducing island-like martensite (hereinafter referred to as MA), which is the hard second phase in the microstructure of bainite single-phase steel pipe, local strain generated around the hard phase in the pipe making stage It was found that the gradient can be relaxed and the decrease in compressive yield stress due to the Bauschinger effect can be suppressed.
しかし、すでに述べたように，優れた脆性き裂伝播停止性能を確保するためには、複相組織化をすることが必然であるので、本発明では複相組織鋼のバウシンガー効果による降伏応力低下の度合いを低減する製造方法について検討し以下の知見を得た。 However, as described above, in order to ensure excellent brittle crack propagation stopping performance, it is necessary to form a multiphase structure. Therefore, in the present invention, the yield stress due to the Bauschinger effect of the multiphase structure steel is used. The manufacturing method for reducing the degree of decrease was examined and the following findings were obtained.
すなわち、
（１）複相組織鋼を構成するフェライト相とベイナイト相との硬度差を低減することで、圧縮降伏応力／引張降伏応力を向上させることができることを見出した。
（２）複相組織鋼の場合でも、Ｍ−Ａなどの硬質第２相を低減することで、圧縮降伏応力／引張降伏応力を向上させることができることを見出した。
That is,
(1) It has been found that the compressive yield stress / tensile yield stress can be improved by reducing the hardness difference between the ferrite phase and the bainite phase constituting the multiphase steel.
(2) It has been found that even in the case of a multiphase steel, the compressive yield stress / tensile yield stress can be improved by reducing the hard second phase such as MA.
（１）については、加速冷却の停止温度を高くすることで所望の組織を得ることが可能であるが、（２）については、冷却停止時に残る未変態オーステナイトの一部が空冷中にＭ−Ａに変態するため、十分な効果が得られなかった。しかしながら、発明者らは、冷却停止温度をより低い温度まで下げて、冷却停止後ただちに急速再加熱を行うことで、ベイナイト相を焼戻し、Ｍ−Ａを分解することによって、バウシンガー効果による降伏応力の低下がより少ないことを知見した。また、冷却停止後急速加熱を行うことは、空冷後炉加熱などにより再加熱するよりも、フェライト相の焼戻しによる集合組織の集積度低下を抑制しながら、ベイナイト相の焼戻し、Ｍ−Ａの分解ができることを知見した。 With regard to (1), it is possible to obtain a desired structure by increasing the stop temperature of accelerated cooling. However, with respect to (2), a part of untransformed austenite remaining at the time of cooling stop is M- Since it transformed to A, a sufficient effect was not obtained. However, the inventors reduced the cooling stop temperature to a lower temperature, and rapidly reheated immediately after stopping the cooling, thereby tempering the bainite phase and decomposing the MA, thereby yield stress due to the Bausinger effect. It was found that there was less decrease in. In addition, rapid heating after cooling is stopped, while tempering of the bainite phase and decomposition of MA are suppressed while suppressing a decrease in the degree of texture accumulation due to tempering of the ferrite phase, rather than reheating by furnace heating after air cooling. I found out that I can do it.
本発明は、上記した知見にさらに検討を加えたもので、
第一の発明は、質量％で、Ｃ：０．０３〜０．０８％、Ｓｉ：０．０１〜０．５０％、Ｍｎ：１．０〜２．０％、Ｐ：０．０１５％以下、Ｓ：０．００５％以下、Ａｌ：０．０８％以下、Ｎｂ：０．００５〜０．０６０％、Ｔｉ：０．００５〜０．０４０％、Ｎ：０．００１〜０．０１０％を含有し、さらに、Ｃｕ：０．１〜０．６％、Ｎｉ：０．１〜１．２％、Ｃｒ：０．０５〜０．４０％、Ｍｏ：０．０５〜０．４０％、Ｖ：０．００５〜０．０７０％の中から選ばれる１種または２種以上を含有し、残部Ｆｅおよび不可避的不純物からなる厚鋼板を管状に曲げ成形し、突合せ部を溶接して鋼管とした後、さらに拡管してなる鋼管であって、当該鋼管の金属組織がフェライト相およびベイナイト相を主体とする複相組織であり、前記フェライト相と前記ベイナイト相との体積分率の合計が８０％以上、残部に含まれる島状マルテンサイト相の体積分率が２％以下であり、前記フェライト相と前記ベイナイト相との平均硬度差が５０以上１５０以下で、Ｘ線回析により得られる管厚中心位置での圧延面の（１００）面の集積度が１．５以上であることを特徴とする圧潰強度に優れた高靱性溶接鋼管である。
The present invention is a further study of the above findings,
1st invention is the mass%, C: 0.03-0.08%, Si: 0.01-0.50%, Mn: 1.0-2.0%, P: 0.015% or less S: 0.005% or less, Al: 0.08% or less, Nb: 0.005-0.060%, Ti: 0.005-0.040%, N: 0.001-0.010% Further, Cu: 0.1 to 0.6%, Ni: 0.1 to 1.2%, Cr: 0.05 to 0.40%, Mo: 0.05 to 0.40%, V : One or more selected from 0.005 to 0.070%, and a thick steel plate composed of the remaining Fe and inevitable impurities is bent into a tubular shape, and the butt portion is welded to form a steel pipe Thereafter, the steel pipe is further expanded, and the metal structure of the steel pipe is a multiphase structure mainly composed of a ferrite phase and a bainite phase, The sum of the volume fractions of the ferrite phase and the bainite phase is 80% or more, the volume fraction of the island martensite phase contained in the balance is 2% or less, and the average hardness difference between the ferrite phase and the bainite phase Is high toughness welding excellent in crushing strength, characterized in that the integration degree of the (100) plane of the rolled surface at the tube thickness center position obtained by X-ray diffraction is 1.5 or more It is a steel pipe.
第二の発明は、さらに、質量％で、Ｃａ：０．０００５〜０．０１００％、Ｍｇ：０．０００５〜０．０１００％、ＲＥＭ：０．０００５〜０．０２００％、Ｚｒ：０．０００５〜０．０３００％の中から選ばれる１種または２種以上を含有することを特徴とする第一の発明に記載の圧潰強度に優れた高靱性溶接鋼管である。 The second invention further includes, in mass%, Ca: 0.0005 to 0.0100%, Mg: 0.0005 to 0.0100%, REM: 0.0005 to 0.0200%, Zr: 0.0005. The high toughness welded steel pipe having excellent crushing strength according to the first invention, characterized by containing one or more selected from -0.0300%.
第三の発明は、第一または第二の発明のいずれかに記載の成分組成を有する鋼を、１０００〜１２００℃に加熱後、９００℃以下の温度域での累積圧下率を５０％以上、二相温度域での累積圧下率を１０〜５０％として、６６０℃以上の温度で熱間圧延を終了後、ただちに冷却速度５〜５０℃／ｓで、２００〜４２０℃まで冷却を行い、冷却停止後、ただちに、４℃／ｓ以上の昇温速度で冷却停止温度より３０℃以上高い温度で、かつ３２０〜５００℃の温度範囲に再加熱した厚鋼板を室温まで冷却後、管状に曲げ成形し、突合せ部を溶接して鋼管とした後、さらに拡管することを特徴とする圧潰強度に優れた高靱性溶接鋼管の製造方法である。 The third invention, after heating the steel having the component composition according to any one of the first or second invention to 1000 to 1200 ° C, the cumulative rolling reduction in the temperature range of 900 ° C or less is 50% or more, Immediately after hot rolling at a temperature of 660 ° C. or higher at a cumulative rolling reduction in the two-phase temperature range of 10 to 50%, cooling is performed to 200 to 420 ° C. at a cooling rate of 5 to 50 ° C./s. Immediately after stopping, the thick steel plate reheated to a temperature range of 320 to 500 ° C. at a temperature higher than the cooling stop temperature by 30 ° C. or more at a temperature rising rate of 4 ° C./s or more is cooled to room temperature, and then bent into a tube shape And it is the manufacturing method of the high toughness welded steel pipe excellent in the crushing strength characterized by welding a butt | matching part and making it a steel pipe, and expanding further.
第四の発明は、前記拡管を拡管率０．５〜１．２５％で行うことを特徴とする第三の発明に記載の圧潰強度に優れた高靱性溶接鋼管の製造方法である。 A fourth invention is a method for producing a high toughness welded steel pipe having excellent crushing strength according to the third invention, wherein the pipe expansion is performed at a pipe expansion ratio of 0.5 to 1.25%.
第五の発明は、前記鋼板を管状に曲げ成形する際に、鋼板に付与される圧縮加工の圧縮率が拡管率の１／４以上であることを特徴とする第三または第四のいずれかに記載の圧潰強度に優れた高靱性溶接鋼管の製造方法である。 According to a fifth invention, in the third or fourth aspect, the compression ratio of the compression processing applied to the steel sheet when the steel sheet is bent into a tubular shape is 1/4 or more of the tube expansion ratio. Is a method for producing a high toughness welded steel pipe having excellent crushing strength.
本発明により、圧潰強度および真円度に優れる石油や天然ガスの輸送とりわけ海底パイプラインに使用させる厚肉高強度ラインパイプ用として好適な厚鋼板を冷間で成形し溶接して製造される高靱性溶接鋼管の製造が可能となり、産業上極めて有効である According to the present invention, a high-temperature steel plate manufactured by cold forming and welding a thick steel plate suitable for transportation of oil and natural gas having excellent crushing strength and roundness, particularly for thick-walled high-strength line pipes used for submarine pipelines. It is possible to manufacture tough welded steel pipes and is extremely effective in industry.
本発明に係る圧潰強度に優れた高靱性溶接鋼管の成分組成、ミクロ組織および管厚中心位置の圧延面の集合組織の形態を説明する。
成分組成
以下に成分組成の限定理由を説明する。なお、成分組成を示す単位は、全て質量％とする。
The composition of the high toughness welded steel pipe excellent in crush strength according to the present invention, the microstructure, and the form of the texture of the rolled surface at the center of the pipe thickness will be described.
Component composition Reasons for limiting the component composition will be described below. In addition, the unit which shows a component composition shall be mass% altogether.
Ｃ：０．０３〜０．０８％
Ｃは焼き入れ性を高め強度確保に重要な元素であるが、０.０３％未満では十分な強度が確保できない。また、０．０８％を超えて添加すると、組織中のマルテンサイトやセメンタイトの体積分率を増加させ、バウシンガー効果を大きくする。よって、Ｃ含有量は、０.０３〜０.０８％の範囲とする。
C: 0.03-0.08%
C is an element that enhances the hardenability and is important for securing the strength, but if it is less than 0.03%, sufficient strength cannot be secured. Moreover, when it adds exceeding 0.08%, the volume fraction of the martensite and cementite in a structure | tissue will be increased, and the Bausinger effect will be enlarged. Therefore, the C content is in the range of 0.03 to 0.08%.
Ｓｉ：０．０１〜０．５０％
Ｓｉは脱酸のため添加するが、０.０１％未満では脱酸効果が十分でなく、０.５％を超えるとマルテンサイト体積分率の増加や溶接性劣化が起こるため、Ｓｉ含有量は０.０１〜０.５０％の範囲とする。さらに好適には、０．０１〜０．２０％の範囲である。
Si: 0.01 to 0.50%
Si is added for deoxidation, but if it is less than 0.01%, the deoxidation effect is not sufficient, and if it exceeds 0.5%, the martensite volume fraction increases and weldability deteriorates. The range is 0.01 to 0.50%. More preferably, it is 0.01 to 0.20% of range.
Ｍｎ：１．０〜２．０％
Ｍｎは強度、靭性向上に有効な元素であるが、１．０％未満ではその効果が十分でなく、２．０％を超えると焼き入れ性が高まりマルテンサイト体積分率の増加、表面硬度の上昇、溶接性劣化を招くため、Ｍｎ含有量は、１．０〜２．０％の範囲とする。
Mn: 1.0-2.0%
Mn is an element effective for improving strength and toughness. However, if it is less than 1.0%, the effect is not sufficient, and if it exceeds 2.0%, the hardenability increases and the martensite volume fraction increases, the surface hardness increases. In order to cause an increase and weldability deterioration, the Mn content is set to a range of 1.0 to 2.0%.
Ｐ：０．０１５％以下
Ｐは不純物元素であり、靭性を劣化させるため、極力低減させることが望ましいが、過度のＰ低減はコストの増大を招くため、Ｐ含有量は０．０１５％以下とする。
P: 0.015% or less P is an impurity element, and it is desirable to reduce it as much as possible in order to degrade toughness. However, excessive P reduction causes an increase in cost, so the P content is 0.015% or less. To do.
Ｓ：０．００５％以下
Ｓは不純物元素であり、靭性を劣化させるため、極力低減させることが望ましいが、過度のＳ低減はコストの増大を招くため、Ｓ含有量は０．００５％以下とする。
S: 0.005% or less S is an impurity element, and it is desirable to reduce as much as possible in order to degrade toughness. However, excessive S reduction causes an increase in cost, so the S content is 0.005% or less. To do.
Ａｌ：０．０８％以下
Ａｌは脱酸剤として添加されるが、０.０８％を超えると鋼の清浄度が低下し、靱性が劣化するため、Ａｌ含有量は０.０８％以下とする。好ましくは、０．０１〜０．０５％の範囲である。
Al: 0.08% or less Al is added as a deoxidizing agent, but if it exceeds 0.08%, the cleanliness of the steel decreases and the toughness deteriorates, so the Al content should be 0.08% or less. . Preferably, it is 0.01 to 0.05% of range.
Ｎｂ：０．００５〜０．０６０％
Ｎｂは制御圧延の効果を高め、組織細粒化により強度、靭性を向上させる元素である。しかし、０.００５％未満では効果がなく、０.０６０％を超えると溶接熱影響部の靭性が劣化するため、Ｎｂ含有量は０.００５〜０.０６０％の範囲とする。
Nb: 0.005 to 0.060%
Nb is an element that enhances the effect of controlled rolling and improves strength and toughness by refining the structure. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.060%, the toughness of the weld heat-affected zone deteriorates.
Ｔｉ：０．００５〜０．０４０％
ＴｉはＴｉＮのピンニング効果により加熱時のオーステナイトの粗大化を抑制し、母材や溶接熱影響部の靭性を改善するために有効な元素である。しかし、０.００５％未満では効果が無く、０.０４０％を超える添加はＴｉＮが粗大化し、逆に溶接熱影響部靭性の劣化を招くため、Ｔｉ含有量は，０.００５〜０.０４０％の範囲とする。さらに、Ｔｉ含有量を０．００５〜０．０２％にすると、より優れた靭性を示す。
Ti: 0.005-0.040%
Ti is an effective element for suppressing the austenite coarsening during heating due to the pinning effect of TiN and improving the toughness of the base metal and the weld heat affected zone. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.040%, TiN becomes coarse and conversely causes deterioration of the weld heat affected zone toughness, so the Ti content is 0.005 to 0.040. % Range. Furthermore, when the Ti content is 0.005 to 0.02%, more excellent toughness is exhibited.
Ｎ：０．００１〜０．０１０％
ＮはＴｉＮのピンニング効果により加熱時のオーステナイトの粗大化を抑制し、母材や溶接熱影響部の靭性を改善するために有効な元素である。しかし、０．００１％以下の含有量では効果がなく、０．０１０％を超えて含有すると添加はＴｉＮの粗大化や固溶Ｎの増大により、逆に溶接熱影響部靱性の劣化を招くため、Ｎの含有量は０．００１〜０．０１０％とする。さらに、Ｎを０．００１〜０．００６％として、質量％の比としてＴｉ／Ｎを１〜５、さらに好ましくは２〜４とすることで、優れた靱性を示す。
N: 0.001 to 0.010%
N is an element effective for suppressing the austenite coarsening during heating by the pinning effect of TiN and improving the toughness of the base metal and the weld heat affected zone. However, if the content is 0.001% or less, there is no effect. If the content exceeds 0.010%, the addition causes the deterioration of the toughness of the weld heat-affected zone due to the coarsening of TiN and the increase in solute N. , N content is 0.001 to 0.010%. Furthermore, excellent toughness is exhibited by setting N to 0.001 to 0.006% and Ti / N to 1 to 5, more preferably 2 to 4 as a ratio by mass.
さらに、鋼板の強度や靱性を向上させるため、以下に示すＣｕ、Ｎｉ、Ｃｒ、Ｍｏ、Ｖの中から選ばれた１種又は２種以上を含有する必要がある。 Furthermore, in order to improve the intensity | strength and toughness of a steel plate, it is necessary to contain the 1 type (s) or 2 or more types chosen from Cu, Ni, Cr, Mo, and V shown below.
Ｃｕ：０．１〜０．６％
Ｃｕは靭性の改善と強度の上昇に有効な元素である。その効果を得るためには、０．１％以上添加することが好ましいが、０．６％を超えて添加すると溶接性の劣化やマルテンサイト体積分率の増加を招くため、Ｃｕを添加する場合はその含有量は０．１〜０．６％の範囲とする。
Cu: 0.1 to 0.6%
Cu is an element effective for improving toughness and increasing strength. In order to obtain the effect, it is preferable to add 0.1% or more, but adding over 0.6% causes deterioration of weldability and an increase in martensite volume fraction, so when adding Cu The content is in the range of 0.1 to 0.6%.
Ｎｉ：０．１〜１．２％
Ｎｉは靭性の改善と強度の上昇に有効な元素である。その効果を得るためには、０．１％以上添加することが好ましいが、１．２％を超えて添加するとコスト的に不利になり、また、溶接熱影響部靱性が劣化するため、Ｎｉを添加する場合はその含有量は０．１〜１．２％の範囲とする。
Ni: 0.1-1.2%
Ni is an element effective for improving toughness and increasing strength. In order to obtain the effect, it is preferable to add 0.1% or more, but adding over 1.2% is disadvantageous in terms of cost, and the weld heat affected zone toughness deteriorates. When added, the content is in the range of 0.1 to 1.2%.
Ｃｒ：０．０５〜０．４０％
ＣｒはＭｎと同様に低Ｃでも十分な強度を得るために有効な元素である。その効果を得るためには、０．０５％以上添加することが好ましいが、０．４０％を超えて添加すると溶接性の劣化やマルテンサイト体積分率の増加を招くため、Ｃｒを添加する場合はその含有量は０．０５〜０．４０％の範囲とする。
Cr: 0.05-0.40%
Cr, like Mn, is an element effective for obtaining sufficient strength even at low C. In order to obtain the effect, it is preferable to add 0.05% or more, but adding over 0.40% causes deterioration of weldability and an increase in martensite volume fraction. Has a content of 0.05 to 0.40%.
Ｍｏ：０．０５〜０．４０％
Ｍｏは焼き入れ性を向上し強度上昇に大きく寄与する元素である。しかし、０.０５％未満ではその効果が得られず、０.４０％を超える添加はマルテンサイト体積分率の増加や溶接熱影響部靭性の劣化を招くため、Ｍｏを添加する場合は、その含有量は０.０５〜０.４０％の範囲とする。さらに好適には０．０５〜０．３％とする。
Mo: 0.05-0.40%
Mo is an element that improves hardenability and greatly contributes to an increase in strength. However, if less than 0.05%, the effect cannot be obtained, and addition exceeding 0.40% leads to an increase in martensite volume fraction and deterioration of weld heat-affected zone toughness. The content is in the range of 0.05 to 0.40%. More preferably, the content is 0.05 to 0.3%.
Ｖ：０．００５〜０．０７０％
Ｖは強度上昇に寄与する元素である。しかし、０.００５％未満では効果がなく、０.０７０％を超えると溶接熱影響部の靭性が劣化するため、Ｖを添加する場合はその含有量は０.００５〜０.０７０％とする。
V: 0.005-0.070%
V is an element contributing to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.070%, the toughness of the weld heat-affected zone deteriorates. Therefore, when V is added, its content is 0.005 to 0.070%. .
さらに、鋼板の欠陥発生の防止や溶接熱影響部の靱性を向上させる場合、以下に示すＣａ、Ｍｇ、ＲＥＭ、Ｚｒの中から選ばれる１種又は２種以上を含有してもよい。 Furthermore, when improving the toughness of a steel plate defect generation | occurrence | production or a welding heat affected zone, you may contain 1 type (s) or 2 or more types chosen from Ca, Mg, REM, and Zr shown below.
Ｃａ：０．０００５〜０．０１００％
ＣａはＭｎＳの形態制御に有効な元素であり、母材靱性の向上に寄与する。その効果を得るためには、０．０００５％以上添加することが好ましいが、０．０１００％を超えて添加するとＣａの酸硫化物が過剰に生成し粗大化やクラスタ状になることにより母材靱性を劣化させることから、Ｃａを添加する場合はその含有量は０．０００５〜０．０１００％の範囲とする。
Ca: 0.0005 to 0.0100%
Ca is an element effective for controlling the morphology of MnS and contributes to improvement of the base material toughness. In order to obtain the effect, it is preferable to add 0.0005% or more. However, if it exceeds 0.0100%, Ca oxysulfide is excessively generated and becomes coarse or clustered to form a base material. In order to deteriorate toughness, when Ca is added, its content is set in the range of 0.0005 to 0.0100%.
Ｍｇ：０．０００５〜０．０１００％
Ｍｇはアルミナクラスタ（Ａｌ２Ｏ３）を、Ａｌ−Ｍｇ系酸化物として微細分散させることで母材靭性向上に寄与する元素である。その効果を得るためには、０．０００５％以上添加することが好ましいが、０．０1００％を越える添加では酸化物の増加により母材靭性の低下が起こるため、Ｍｇを添加する場合はその含有量は０．０００５〜０．０１００％の範囲とする。
Mg: 0.0005 to 0.0100%
Mg is an element that contributes to improving the toughness of the base material by finely dispersing alumina clusters (Al 2 O 3 ) as an Al—Mg-based oxide. In order to obtain the effect, it is preferable to add 0.0005% or more. However, if the addition exceeds 0.0100%, the toughness of the base metal decreases due to an increase in oxide. The amount is in the range of 0.0005 to 0.0100%.
ＲＥＭ：０．０００５〜０．０２００％
ＲＥＭ（Rare Earth Metals：希土類金属）はＣａと同様、ＭｎＳの形態制御に有効な元素であり、母材靭性の向上に寄与する。その効果を得るためには、０．０００５％以上添加することが好ましいが、０．０２００％以上の添加は、ＲＥＭの酸硫化物が過剰に生成し、母材靭性を劣化させるため、ＲＥＭを添加する場合はその含有量は０．０００５〜０．０２００％の範囲とする。
REM: 0.0005 to 0.0200%
REM (Rare Earth Metals), like Ca, is an element that is effective in controlling the morphology of MnS, and contributes to the improvement of base metal toughness. In order to obtain the effect, it is preferable to add 0.0005% or more. However, addition of 0.0200% or more causes excessive generation of REM oxysulfide and deteriorates the base material toughness. When adding, the content shall be 0.0005 to 0.0200% of range.
Ｚｒ：０．０００５〜０．０３００％
ＺｒはＣａと同様、ＭｎＳの形態制御に有効な元素であり、母材靭性の向上に寄与する。その効果を得るためには、０．０００５％以上添加することが好ましいが、０．０３００％超えての添加は、Ｚｒの酸硫化物が過剰に生成し、母材靭性を劣化させ、さらにＴｉＮと複合化することにより溶接熱影響部靱性を劣化させるため、Ｚｒを添加する場合はその含有量は０．０００５〜０．０３００％の範囲とする。
Zr: 0.0005 to 0.0300%
Zr, like Ca, is an element effective for controlling the morphology of MnS and contributes to the improvement of the base metal toughness. In order to obtain the effect, it is preferable to add 0.0005% or more. However, if it exceeds 0.0300%, Zr oxysulfide is excessively generated, the base material toughness is deteriorated, and TiN is further added. Therefore, when Zr is added, the content is made 0.0005 to 0.0300%.
上記以外の残部はＦｅおよび不可避的不純物とする。
なお、Ｂを含有することにより熱間圧延中のフェライト相の生成が抑制され、フェライト相の加工による集合組織の発達が困難になるため、本発明ではＢは不可避的不純物として取り扱い、好ましくは、０．０００５％以下とする。
The balance other than the above is Fe and inevitable impurities.
In addition, since the formation of the ferrite phase during hot rolling is suppressed by containing B, and the development of the texture by the processing of the ferrite phase becomes difficult, in the present invention, B is treated as an inevitable impurity, 0.0005% or less.
ミクロ組織
本発明では、金属組織の形態および体積分率を規定する。金属組織はフェライト相とベイナイト相を主体とする。フェライト相は圧延中に加工することにより集合組織を発達させ，脆性き裂伝播停止性能を向上させるために必須の組織である。一方、強度を確保するためにはベイナイト相やマルテンサイト相などの硬質相を導入する必要があるが、マルテンサイト相ではフェライト相との硬度差を所望の範囲にすることができず、バウシンガー効果による降伏応力の低下を十分に抑制できず，優れた圧縮強度を得ることができないため、フェライト相とベイナイト相の合計体積分率を８０％以上とする。残部は、Ｍ−Ａ、マルテンサイト、パーライト、セメンタイトなどであるが、これらは、できるだけ少ないことが好ましい。
Microstructure In the present invention, the form and volume fraction of the metal structure are defined. The metal structure is mainly composed of a ferrite phase and a bainite phase. The ferrite phase is an indispensable structure for developing the texture by processing during rolling and improving the brittle crack propagation stopping performance. On the other hand, in order to ensure strength, it is necessary to introduce a hard phase such as a bainite phase or a martensite phase. However, in the martensite phase, the hardness difference from the ferrite phase cannot be brought into a desired range, and the bausinger Since the decrease in yield stress due to the effect cannot be sufficiently suppressed and excellent compressive strength cannot be obtained, the total volume fraction of the ferrite phase and the bainite phase is set to 80% or more. The balance is MA, martensite, pearlite, cementite, etc., but these are preferably as small as possible.
なかでも、硬質第２相と母相の周辺に発生する局所的なひずみ勾配による背応力の発生を防止し、バウシンガー効果による圧縮降伏応力低下を抑制するため，金属組織中においてＭ−Ａの体積分率を２％以下とする。 In particular, in order to prevent the occurrence of back stress due to local strain gradient generated around the hard second phase and the parent phase, and to suppress the decrease in compressive yield stress due to the Bauschinger effect, The volume fraction is 2% or less.
その他残存組織として、二相域から加速冷却を開始した場合に，数％程度のパーライトが観察されるほか、Ｍ−Ａの分解生成物としてセメンタイトが観察される。また、後述のように圧延終了後の冷却速度が過大であるとマルテンサイトが混入する。 As the remaining structure, when accelerated cooling is started from the two-phase region, about several percent of pearlite is observed, and cementite is observed as a decomposition product of MA. Further, as will be described later, martensite is mixed when the cooling rate after the end of rolling is excessive.
ミクロ組織間の硬度差：５０〜１５０
本発明では、主要な金属組織として規定したフェライト相とベイナイト相との硬度差を規定する。荷重負荷時の金属組織内の局所的なひずみ勾配は、先に述べた硬質第２相と母相の間だけでなく、母相である複相組織鋼の軟質相と硬質相の間にも発生し、バウシンガー効果による降伏応力低下を助長するため、鋼管の圧縮応力が低下する。
Hardness difference between microstructures: 50-150
In the present invention, the hardness difference between the ferrite phase and the bainite phase defined as the main metal structure is defined. The local strain gradient in the metal structure at the time of loading is not only between the hard second phase and the parent phase described above, but also between the soft phase and the hard phase of the multiphase steel that is the parent phase. It is generated and promotes a decrease in yield stress due to the Bauschinger effect, so that the compressive stress of the steel pipe is reduced.
したがって、母相の軟質相であるフェライト相と母相の硬質相であるベイナイト相の平均硬度差を１５０以下にすることで、バウシンガー効果による降伏応力低下を抑制することができる。一方、強度確保の観点およびセパレーションの発生を容易にするためには、下限を５０にするものとする。 Therefore, by reducing the average hardness difference between the ferrite phase, which is the soft phase of the parent phase, and the bainite phase, which is the hard phase of the parent phase, to 150 or less, it is possible to suppress a decrease in yield stress due to the Bauschinger effect. On the other hand, the lower limit is set to 50 in order to ensure strength and facilitate the generation of separation.
なお、平均硬さの測定方法については、荷重０．９８Ｎ以下のマイクロビッカース試験機により任意の２０点以上を測定し、その平均値をとることが好ましい。 In addition, about the measuring method of average hardness, it is preferable to measure 20 or more arbitrary points with a micro Vickers tester with a load of 0.98 N or less, and take the average value.
板厚中心位置における集合組織
優れた脆性き裂伝播停止性能を得るためには、脆性き裂発生時にいわゆるき裂進展のセパレーションを発生させることが必要である。セパレーションは、一般的に圧延面に（１００）面と（１１１）面が発達している際に発生しやすくなることが知られている。本発明では、脆性き裂伝播停止性能の評価法としてＤＷＴＴ試験を採用し、様々な鋼板についてＤＷＴＴ試験を行った。
Texture at the center of the plate thickness In order to obtain excellent brittle crack propagation stopping performance, it is necessary to generate a so-called separation of crack growth when a brittle crack is generated. It is known that separation tends to occur when a (100) plane and a (111) plane are generally developed on a rolled surface. In the present invention, the DWTT test was adopted as a method for evaluating the brittle crack propagation stopping performance, and various steel sheets were subjected to the DWTT test.
その結果、ＤＷＴＴ試験の延性破面率とＸ線回析により得られる管厚中心位置での圧延面の（１００）面の集積度とがよい相関があることが判明し、フェライト相とベイナイト相の硬度差が５０以上の場合、（１００）面の集積度を１．５以上とすることで、本発明範囲の溶接鋼管で優れた脆性き裂伝播停止性能が得られることがわかったので、（１００）面の集積度の下限を１．５とした。 As a result, it was found that there is a good correlation between the ductile fracture surface ratio of the DWTT test and the degree of integration of the (100) plane of the rolled surface at the tube thickness center position obtained by X-ray diffraction, and the ferrite phase and the bainite phase When the hardness difference is 50 or more, it was found that by setting the degree of integration of the (100) plane to 1.5 or more, excellent brittle crack propagation stopping performance can be obtained with the welded steel pipe in the range of the present invention. The lower limit of the degree of integration of the (100) plane was 1.5.
なお、ここで（１００）面の集積度とは、集合組織のないランダムな標準試料における（２００）面からのＸ線回折強度に対する、管厚中心位置から鋼板圧延面に平行に採取した板面における（２００）面からのＸ線回析強度の比をいう。 Here, the degree of integration of the (100) plane is a plate plane taken in parallel with the steel sheet rolling plane from the tube thickness center position with respect to the X-ray diffraction intensity from the (200) plane in a random standard sample without a texture. The ratio of the X-ray diffraction intensity from the (200) plane.
次に、本発明に係る鋼管の素材とする厚鋼板の好適な製造方法について説明する。製造方法においては、スラブ加熱温度、熱間圧延、加速冷却、および加速冷却後の再加熱条件を規定する。
加熱温度、圧延終了温度、冷却停止温度、再加熱温度で規定している温度は鋼板全体の平均温度とする。平均温度は、スラブもしくは鋼板の表面温度より、板厚、熱伝導率等のパラメータを考慮して、計算により求めたものである。
また、冷却速度は、冷却開始温度と冷却停止温度（４００〜６００℃）との温度差をその冷却を行うのに要した時間で割った平均冷却速度とする。
Next, the suitable manufacturing method of the thick steel plate used as the raw material of the steel pipe concerning this invention is demonstrated. In the manufacturing method, slab heating temperature, hot rolling, accelerated cooling, and reheating conditions after accelerated cooling are defined.
The temperature defined by the heating temperature, rolling end temperature, cooling stop temperature, and reheating temperature is the average temperature of the entire steel sheet. The average temperature is obtained by calculation based on the surface temperature of the slab or steel plate, taking into account parameters such as plate thickness and thermal conductivity.
The cooling rate is an average cooling rate obtained by dividing the temperature difference between the cooling start temperature and the cooling stop temperature (400 to 600 ° C.) by the time required for the cooling.
スラブ加熱温度：１０００〜１２００℃
スラブをオーステナイト化しつつ、最低限のＮｂの固溶量を得るため、下限温度は１０００℃である。一方、１２００℃を超える温度までスラブを加熱すると、ＴｉＮによるピンニング効果が弱まり、オーステナイト粒が著しく成長し、母材靭性が劣化する。このため、スラブ加熱温度は１０００〜１２００℃の範囲とする。
Slab heating temperature: 1000-1200 ° C
The minimum temperature is 1000 ° C. in order to obtain a minimum amount of Nb solid solution while austenizing the slab. On the other hand, when the slab is heated to a temperature exceeding 1200 ° C., the pinning effect due to TiN is weakened, austenite grains grow significantly, and the base material toughness deteriorates. For this reason, slab heating temperature shall be 1000-1200 degreeC.
９００℃以下の温度域での累積圧下率：５０％以上
本発明に係る鋼では、Ｎｂ添加によって９００℃以下はオーステナイト未再結晶温度領域である。この温度域以下において累積で大圧下の圧延を行うことにより、オーステナイト粒を伸展させ、特に板厚方向で細粒とし母材靭性を向上させる。累積圧下率が５０％未満の場合は、細粒化が十分でなく靱性が劣化するため、９００℃以下の温度域での累積圧下率は５０％以上とする。
Cumulative rolling reduction in a temperature range of 900 ° C. or less: 50% or more In the steel according to the present invention, 900 ° C. or less is an austenite non-recrystallization temperature region due to Nb addition. Austenite grains are extended by rolling under a large cumulative pressure below this temperature range, and in particular, the base material toughness is improved by making fine grains in the plate thickness direction. When the cumulative rolling reduction is less than 50%, the fine reduction is not sufficient and the toughness deteriorates, so the cumulative rolling reduction in the temperature range of 900 ° C. or lower is set to 50% or more.
二相温度域での累積圧下率：１０〜５０％
Ａｒ３点〜Ａｒ１点のフェライト−オーステナイト二相温度域で熱間圧延を行うことによってオーステナイト未再結晶域圧延で細粒化したオーステナイトをさらに微細化する。
Cumulative rolling reduction in two-phase temperature range: 10-50%
Austenite refined by austenite non-recrystallization zone rolling is further refined by performing hot rolling in the ferrite-austenite two-phase temperature range of Ar 3 to Ar 1 .
さらに、フェライトに加工を加えることによってフェライト強化による高強度化とＤＷＴＴなどの脆性き裂伝播停止性能評価試験で、試験片の破面にセパレーションを発生させるのに必要な集合組織形態を実現し、優れた脆性き裂伝播停止性能とすることが可能となる。 In addition, by applying processing to ferrite, strengthening by ferrite strengthening and brittle crack propagation stopping performance evaluation test such as DWTT, realize the texture form necessary to generate separation on the fracture surface of the test piece, It is possible to achieve excellent brittle crack propagation stopping performance.
二相温度域の累積圧下量が１０％未満では、集合組織の発達が少なくセパレーションの発生が十分でなく脆性き裂伝播停止特性の向上が得られない。一方、累積圧下率が５０％を超えると、フェライトへの過剰な加工によりフェライトが脆化し、母材靭性が劣化する。このため、二相温度域での累積圧下率を１０〜５０％の範囲とする。 If the cumulative reduction in the two-phase temperature range is less than 10%, the texture development is small and the occurrence of separation is not sufficient, and the improvement of brittle crack propagation stopping characteristics cannot be obtained. On the other hand, if the cumulative rolling reduction exceeds 50%, the ferrite becomes brittle due to excessive processing to ferrite, and the base metal toughness deteriorates. For this reason, the cumulative rolling reduction in the two-phase temperature range is set to a range of 10 to 50%.
圧延終了温度：６６０℃以上
圧延終了温度が６６０℃未満の場合、フェライト変態が進行して加速冷却の効果が小さくなり、かつフェライトが粗大化することにより母材靭性が劣化するため、圧延終了温度は６６０℃以上とする。
Rolling end temperature: 660 ° C. or higher When the rolling end temperature is lower than 660 ° C., the ferrite transformation proceeds and the effect of accelerated cooling is reduced, and the base metal toughness is deteriorated due to the coarsening of the ferrite. Is 660 ° C. or higher.
冷却速度：５〜５０℃／ｓ
圧延終了後に生成するフェライトは加工されていないため、強度、靭性確保の観点からは有害である。したがって、圧延終了後ただちに５℃／ｓ以上の冷却速度で加速冷却を行い、未変態オーステナイトをベイナイト組織に変態させてフェライトの発生を防止し、母材靭性を損なわずに強度を向上させる。一方で、本発明のように冷却停止温度を低くする場合は、冷却速度が過剰であるとベイナイト組織の中にマルテンサイト組織が混入する。５０℃／ｓを超える冷却速度の場合その傾向が顕著であり所望の組織形態が得られないため、上限を５０℃／ｓとする。
Cooling rate: 5-50 ° C./s
Since ferrite produced after rolling is not processed, it is harmful from the viewpoint of securing strength and toughness. Therefore, immediately after the rolling is completed, accelerated cooling is performed at a cooling rate of 5 ° C./s or more, and untransformed austenite is transformed into a bainite structure to prevent the generation of ferrite, and the strength is improved without impairing the base material toughness. On the other hand, when the cooling stop temperature is lowered as in the present invention, if the cooling rate is excessive, the martensite structure is mixed in the bainite structure. In the case of a cooling rate exceeding 50 ° C./s, the tendency is remarkable and a desired tissue form cannot be obtained. Therefore, the upper limit is set to 50 ° C./s.
冷却停止温度：２００〜４２０℃
再加熱後の引張強さを６００ＭＰａ以上とするため、冷却停止温度を４２０℃以下として、鋼板の加速冷却前に未変態オーステナイトであった部分をベイナイト組織とする。冷却停止温度が４２０℃を超えると変態温度が高く、十分に鋼板を高強度化できないため、上限を４２０℃とする。また、冷却停止温度が２００℃を下回ると、マルテンサイトの混入が避けられないことおよび鋼管素材の板内ひずみが大きく鋼管製造時に矯正のために過大なプレス圧縮もしくは拡管を与える必要があるため、下限を２００℃とする。冷却方法については製造プロセスによって任意の冷却設備を用いることが可能であり、例えば水冷方式の加速冷却設備が利用できる。
Cooling stop temperature: 200-420 ° C
In order to set the tensile strength after reheating to 600 MPa or more, the cooling stop temperature is set to 420 ° C. or lower, and the portion that was untransformed austenite before the accelerated cooling of the steel sheet is made a bainite structure. If the cooling stop temperature exceeds 420 ° C, the transformation temperature is high and the steel sheet cannot be sufficiently strengthened, so the upper limit is set to 420 ° C. Also, if the cooling stop temperature is below 200 ° C., it is unavoidable that martensite is mixed, and because the strain in the plate of the steel pipe material is large, it is necessary to give excessive press compression or expansion for correction at the time of steel pipe production. The lower limit is 200 ° C. As a cooling method, any cooling equipment can be used depending on the manufacturing process, and for example, a water-cooled accelerated cooling equipment can be used.
再加熱処理：３２０〜５００℃
本発明において、再加熱処理は、重要な熱処理で、複相組織を有する加速冷却ままの鋼板のベイナイト組織を焼き戻してフェライト相とベイナイト相との硬度差を低減し、またＭ−Ａを分解するために行う。
Reheating treatment: 320 to 500 ° C
In the present invention, the reheating treatment is an important heat treatment, tempering the bainite structure of the steel sheet having accelerated cooling and having a multiphase structure to reduce the hardness difference between the ferrite phase and the bainite phase, and decompose MA. To do.
フェライト相とベイナイト相の硬度差の低減とＭ−Ａ分解を達成するためには、３２０℃以上に再加熱する必要がある。また、冷却停止温度よりも３０℃以上の温度に昇温しなければ、再加熱の効果が得られないため、下限を冷却停止温度よりも３０℃以上高くなおかつ３２０℃以上の温度とする。 In order to reduce the hardness difference between the ferrite phase and the bainite phase and achieve the MA decomposition, it is necessary to reheat to 320 ° C. or higher. Further, since the effect of reheating cannot be obtained unless the temperature is raised to 30 ° C. or higher than the cooling stop temperature, the lower limit is set to 30 ° C. or higher and 320 ° C. or higher than the cooling stop temperature.
一方、５００℃以上に加熱すると、焼戻し効果が顕著となり引張強度が著しく低下するだけでなく、フェライト相とベイナイト相の硬度差が小さくなりすぎることと、集合組織の集積度が低下することとの重畳でセパレーションの発生量が低下する現象が起こり、脆性き裂伝播停止性能が低下するため、上限を５００℃とする。 On the other hand, when heated to 500 ° C. or higher, not only the tempering effect is remarkable and the tensile strength is remarkably lowered, but also the hardness difference between the ferrite phase and the bainite phase is too small, and the accumulation degree of the texture is lowered. Since the phenomenon that the generation amount of separation decreases due to superposition occurs and the brittle crack propagation stopping performance decreases, the upper limit is set to 500 ° C.
再加熱処理時の昇温速度：４℃／ｓ以上
冷却停止後急速加熱を行うことは、空冷後炉加熱などにより再加熱するよりも、フェライト相の焼戻しによる集合組織の集積度低下を抑制しながら、ベイナイト相の焼戻し、Ｍ−Ａの分解ができ、生産性の観点からみても有利であるため、冷却停止後直ちに、急速加熱により、４℃／ｓ以上、望ましくは６℃／ｓ以上の昇温速度で再加熱するものとする。再加熱後の冷却過程は特に規定しないが、空冷とするとＭ−Ａの再生成を防止できるため好適である。
Heating rate during reheating: 4 ° C / s or more Rapid heating after stopping cooling suppresses a decrease in texture accumulation due to tempering of the ferrite phase, rather than reheating by furnace heating after air cooling. However, the bainite phase can be tempered and the MA can be decomposed, which is advantageous from the viewpoint of productivity. Immediately after the cooling is stopped, it is rapidly heated to 4 ° C./s or more, preferably 6 ° C./s or more. It shall be reheated at a rate of temperature increase. The cooling process after reheating is not particularly defined, but air cooling is preferable because it can prevent the regeneration of MA.
加速冷却後の再加熱を行うための設備として、冷却設備の下流側に加熱装置を設置する。加熱装置としては、鋼板表面と板厚中央部で温度差を発生させることが容易な誘導加熱装置を用いる事が好ましい。 As equipment for performing reheating after accelerated cooling, a heating device is installed on the downstream side of the cooling equipment. As the heating device, it is preferable to use an induction heating device that can easily generate a temperature difference between the surface of the steel plate and the central portion of the plate thickness.
上述した製造方法を実施する設備として、たとえば圧延ラインの上流から下流側に向かって熱間圧延機、冷却装置、誘導加熱装置、ホットレベラーを逐次配置したものが好適である。
誘導加熱装置あるいは他の熱処理装置を、圧延設備である熱間圧延機およびその出側に配置される冷却装置と同一ライン上に設置する事によって、圧延、加速冷却終了後迅速に再加熱処理が行えるので、加速冷却後の鋼板温度を過度に低下させることなく加熱することが可能である。
続いて、本発明に係る鋼管の好適な製造方法について説明する。本発明において好適な鋼管の製造方法は、常法を基本とするものの、拡管率および冷間曲げ時に付与される圧縮加工の圧縮率を規定する。
As equipment for performing the above-described manufacturing method, for example, a hot rolling mill, a cooling device, an induction heating device, and a hot leveler are sequentially arranged from the upstream side to the downstream side of the rolling line.
By installing an induction heating device or other heat treatment device on the same line as the hot rolling mill that is the rolling equipment and the cooling device arranged on the outlet side, the reheating treatment can be performed quickly after the completion of rolling and accelerated cooling. Since it can be performed, it is possible to heat the steel sheet after accelerated cooling without excessively reducing the steel sheet temperature.
Then, the suitable manufacturing method of the steel pipe concerning this invention is demonstrated. In the present invention, a steel pipe production method suitable for the present invention is based on a conventional method, but defines the pipe expansion ratio and the compression ratio of the compression process applied during cold bending.
拡管率
一般的に厚肉の高強度ＵＯＥ鋼管は、０．９〜１．２％程度の範囲の拡管率で造管を行う。一方で、鋼管周方向の圧縮強度を確保するために、拡管率を低減することが効果的であり、圧縮強度を確保する必要のあるＵＯＥ鋼管は、通常の範囲の下限もしくは、それよりも小さい拡管率の範囲（例えば０．６〜１．０％）で造管される。
Tube expansion rate Generally, a thick high-strength UOE steel pipe is piped at a tube expansion rate of about 0.9 to 1.2%. On the other hand, in order to secure the compressive strength in the circumferential direction of the steel pipe, it is effective to reduce the pipe expansion rate, and the UOE steel pipe that needs to secure the compressive strength is the lower limit of the normal range or smaller than that. The pipe is formed in a range of the expansion ratio (for example, 0.6 to 1.0%).
本発明では、真円度の確保の観点から下限値を０．５％とするが、母材の製造方法により圧縮強度を向上させることが可能であるので、従来の手法よりも高い拡管率（例えば１．１％以上）での製造が可能である。本発明の鋼板を用いることにより、１．２５％までの拡管率で所望の圧縮強度特性が得られる。 In the present invention, the lower limit is set to 0.5% from the viewpoint of securing roundness, but the compression strength can be improved by the manufacturing method of the base material, so that the tube expansion rate ( For example, it can be manufactured at 1.1% or more. By using the steel sheet of the present invention, desired compression strength characteristics can be obtained with a tube expansion rate of up to 1.25%.
冷間曲げ時に付与される圧縮加工の圧縮率
一般的に、厚肉の高強度ＵＯＥ鋼管は、０．３〜０．５％程度（圧縮率/拡管率が１／３以上）の範囲の圧縮率の圧縮加工が付与される冷間曲げにより造管を行う。鋼管周方向の圧縮強度を確保するために、冷間曲げ時に鋼板に付与される圧縮加工の圧縮率を高くすることが効果的である。
Compression rate of compression applied during cold bending Generally, thick high-strength UOE steel pipes are compressed in the range of about 0.3 to 0.5% (compression ratio / expansion ratio is 1/3 or more). Pipe making is performed by cold bending to which a compression process is applied. In order to ensure the compressive strength in the circumferential direction of the steel pipe, it is effective to increase the compression rate of the compression process applied to the steel sheet during cold bending.
本発明では、母材、すなわち鋼管素材である厚鋼板の圧縮強度が向上しているので、従来の手法よりも低い圧縮率での製造が可能であり、拡管率の１／４以上の圧縮率であれば所望の圧縮強度特性が得られる。 In the present invention, since the compressive strength of the base material, that is, the thick steel plate, which is a steel pipe material, is improved, it is possible to manufacture at a lower compression ratio than the conventional method, and the compression ratio is 1/4 or more of the pipe expansion ratio. If so, desired compression strength characteristics can be obtained.
表１に示す化学成分の鋼（鋼種Ａ〜Ｈ）を連続鋳造法によりスラブとし、加熱したスラブを熱間圧延により圧延した後、ただちに水冷型の冷却設備を用いて加速冷却を行い、誘導加熱装置を用いて再加熱を行って板厚８ｍｍおよび２６ｍｍの厚鋼板（Ｎｏ．１〜１９）を製造した。誘導加熱装置は、冷却設備と同一ライン上に設置した。一部、比較のため、誘導加熱装置でなく、一般的な熱処理炉（雰囲気炉）を用いて再加熱を行った。 Steel of the chemical composition shown in Table 1 (steel types A to H) is made into a slab by a continuous casting method, and the heated slab is rolled by hot rolling, and then immediately accelerated and cooled using a water-cooled cooling facility, and induction heating is performed. Reheating was performed using an apparatus to produce thick steel plates (No. 1 to 19) having a thickness of 8 mm and 26 mm. The induction heating device was installed on the same line as the cooling facility. For comparison, reheating was performed using a general heat treatment furnace (atmosphere furnace) instead of an induction heating apparatus.
各鋼板（Ｎｏ．１〜１９）の製造条件を表２に示す。 Table 2 shows the production conditions of each steel plate (No. 1 to 19).
なお、加熱温度、圧延終了温度、冷却開始および停止温度、再加熱温度は鋼板全体の平均温度とした。平均温度は、スラブもしくは鋼板の表面温度より、板厚、熱伝導率等のパラメータから計算により求めた。 The heating temperature, rolling end temperature, cooling start and stop temperature, and reheating temperature were the average temperature of the entire steel sheet. The average temperature was obtained from the surface temperature of the slab or steel plate by calculation from parameters such as plate thickness and thermal conductivity.
加速冷却速度は、加速冷却開始温度と加速冷却停止温度との温度差をその冷却を行うのに要した時間で割った平均冷却速度とした。 The accelerated cooling rate was an average cooling rate obtained by dividing the temperature difference between the accelerated cooling start temperature and the accelerated cooling stop temperature by the time required for the cooling.
鋼管のミクロ組織の分率は、管厚中心位置で４００倍で組織観察した１０枚の光学顕微鏡写真の画像解析からフェライト相とベイナイト相の合計の面積分率を平均して求め、鋼管中に均一にそれらの組織が分散していると仮定して、前記面積分率の値が体積分率の値に等しいものとみなした。 The fraction of the microstructure of the steel pipe is obtained by averaging the total area fraction of the ferrite phase and the bainite phase from image analysis of 10 optical micrographs observed at 400 times the center of the tube thickness. Assuming that the tissues were uniformly distributed, the area fraction value was considered equal to the volume fraction value.
同様に、鋼管のＭ−Ａ体積分率は、管厚中心位置で２０００倍で組織観察した５枚のＳＥＭ（走査型電子顕微鏡）写真の画像解析から面積分率を平均して求め、鋼管中に均一に第２相が分散していると仮定して、前記面積分率の値が体積分率の値に等しいものとみなした。 Similarly, the MA volume fraction of the steel pipe is obtained by averaging the area fraction from image analysis of five SEM (scanning electron microscope) photographs observed at 2000 times the structure at the center of the pipe thickness. Assuming that the second phase is uniformly distributed, the area fraction value is considered to be equal to the volume fraction value.
フェライト相とベイナイト相の硬度差は、荷重０．９８Ｎのマイクロビッカース試験機により板厚１／４ｔ位置で各相それぞれ４０点以上を測定し、各相の硬度の平均値の差を求めることで得た。 The hardness difference between the ferrite phase and the bainite phase is obtained by measuring 40 points or more of each phase at a thickness of 1/4 t with a micro Vickers tester with a load of 0.98 N, and calculating the difference in the average value of the hardness of each phase. Obtained.
管厚中心位置における圧延面の（１００）面の集積度は、集合組織のないランダムな標準試料における（２００）面からのＸ線回折強度に対する、管厚中心位置から圧延面に平行に採取した板面における（２００）面からのＸ線回析強度の比を用いた。 The degree of integration of the (100) plane of the rolled surface at the tube thickness center position was collected parallel to the rolled surface from the tube thickness center position with respect to the X-ray diffraction intensity from the (200) plane in a random standard sample without texture. The ratio of the X-ray diffraction intensity from the (200) plane on the plate surface was used.
鋼管周方向の引張降伏応力は、シーム溶接部から周方向に９０°もしくは２７０°の管体から全厚サンプルを採取し、プレス機で平滑化した後、全厚試験片を作製した。引張降伏応力は、各条件で２本ずつ測定し平均値で評価した。 Regarding the tensile yield stress in the circumferential direction of the steel pipe, a full-thickness specimen was prepared from a full-thickness sample taken from a tubular body of 90 ° or 270 ° in the circumferential direction from the seam weld and smoothed by a press. Two tensile yield stresses were measured under each condition and evaluated by an average value.
鋼管周方向の圧縮降伏応力は、シーム溶接部から周方向に９０°もしくは２７０°の管体の内表面から１ｍｍ以上管厚中心側からＡＳＴＭ Ｅ９準拠の直径２０ｍｍ、長さ６０ｍｍの丸棒試験片を採取し、各条件で２本ずつ測定した０．２％耐力の平均値で圧縮降伏応力を評価した。 Compressive yield stress in the circumferential direction of the steel pipe is 1 mm or more from the inner surface of the pipe body at 90 ° or 270 ° in the circumferential direction from the seam welded portion, and a round bar test piece having a diameter of 20 mm and a length of 60 mm in accordance with ASTM E9. The compressive yield stress was evaluated based on an average value of 0.2% proof stress measured two by two under each condition.
脆性き裂伝播停止特性はＤＷＴＴ試験で評価した。ＤＷＴＴ試験の延性破面率は、管厚１／２ｔ位置から採取した１９ｍｍに減厚したＤＷＴＴ試験片（管厚８ｍｍの鋼管は全厚）を−４７℃で各２本ずつ行い、延性破面率の平均を求めた。延性破面率は、８５％以上を本発明で必要な値とした。 The brittle crack propagation stopping property was evaluated by the DWTT test. The ductile fracture surface ratio of the DWTT test was determined by measuring two DWTT specimens (total thickness of steel pipes with a pipe thickness of 8 mm) sampled from the pipe thickness of 1 / 2t at 19 mm at -47 ° C. The average of the rate was obtained. The ductile fracture surface ratio was set to 85% or more as a necessary value in the present invention.
表３に得られた試験結果を示す。 Table 3 shows the test results obtained.
鋼管Ｎｏ．１、２、１２、１３、１４はいずれも本発明の成分範囲、組織形態範囲、製造条件範囲を満たすため、所望の引張強度特性、圧縮降伏応力特性、DWTT特性が得られている。一方、その他の鋼管では、本発明の範囲外であるため、これらいずれかの特性を満たしていない。 Steel pipe No. Since 1, 2, 12, 13, and 14 all satisfy the component range, structure morphology range, and production condition range of the present invention, desired tensile strength characteristics, compressive yield stress characteristics, and DWTT characteristics are obtained. On the other hand, other steel pipes are outside the scope of the present invention, and therefore do not satisfy any of these characteristics.
鋼管Ｎｏ．３および４は、拡管率、あるいは、圧縮率／拡管率の比が本発明の範囲外であるため、圧縮降伏応力／引張降伏応力の比が小さくなっている。 Steel pipe No. 3 and 4 have a ratio of compressive yield stress / tensile yield stress that is small because the ratio of tube expansion or the ratio of compressibility / expansion ratio is outside the scope of the present invention.
鋼管Ｎｏ．５は、加熱温度が高いため、圧延前組織の粗大化が圧延後も受け継がれ、靱性が劣化し、ＤＷＴＴ特性を満足していない。鋼管Ｎｏ．６は、圧延終了温度が高くベイナイト単相組織となっているため、集合組織が発達しておらず、ＤＷＴＴ特性が劣化している。鋼管Ｎｏ．７は、二相域での圧延を行っていないため、フェライトを加工していないことによる強度不足ならびに集合組織が発達していないことによるＤＷＴＴ特性の劣化がみられる。 Steel pipe No. In No. 5, since the heating temperature is high, the coarsening of the structure before rolling is inherited even after rolling, the toughness is deteriorated, and the DWTT characteristics are not satisfied. Steel pipe No. No. 6 has a high rolling end temperature and has a bainite single-phase structure, so that the texture is not developed and the DWTT characteristics are deteriorated. Steel pipe No. In No. 7, since rolling in the two-phase region is not performed, the strength is insufficient due to the fact that the ferrite is not processed, and the DWTT characteristic is deteriorated due to the fact that the texture is not developed.
鋼管Ｎｏ．８は、冷却速度が大きすぎるために、加速冷却前に未変態オーステナイトであった部分がベイナイト相とマルテンサイト相の混合組織に変態し、フェライト相とこの混合組織（ベイナイト相＋マルテンサイト相）との硬度差が大きくなりすぎて、圧縮降伏応力／引張降伏応力が小さくなっている。鋼管Ｎｏ．９は、冷却停止温度および再加熱温度が高すぎるため、強度の低下とセメンタイトの粗大化によるＤＷＴＴ特性の劣化が見られる。 Steel pipe No. No. 8, because the cooling rate is too high, the part that was untransformed austenite before accelerated cooling was transformed into a mixed structure of bainite phase and martensite phase, and the ferrite phase and this mixed structure (bainite phase + martensite phase) And the hardness difference is too large, and the compressive yield stress / tensile yield stress is small. Steel pipe No. In No. 9, since the cooling stop temperature and the reheating temperature are too high, a decrease in strength and a deterioration in DWTT characteristics due to coarsening of cementite are observed.
鋼管Ｎｏ．１０は、冷却停止後の再加熱を行っていないため、フェライト相とベイナイト相の硬度差が大きいことおよびＭ−Ａ体積分率が大きいことによる、圧縮降伏応力／引張降伏応力の低下がみられる。鋼管Ｎｏ．１１は、冷却停止温度が高く、再加熱を行っていないため、Ｍ−Ａ体積分率が大きいことによる圧縮降伏応力／引張降伏応力の低下が見られる。 Steel pipe No. No. 10 shows no reduction in compressive / stress yield stress due to a large hardness difference between the ferrite phase and the bainite phase and a large MA volume fraction because no reheating was performed after cooling was stopped. . Steel pipe No. No. 11 has a high cooling stop temperature and no reheating, and therefore a decrease in compressive yield stress / tensile yield stress due to a large MA volume fraction is observed.
鋼管Ｎｏ．１５、１６は、それぞれＣ、Ｍｎ、Ｎｂが本発明の請求範囲よりも大きいため、Ｍ−Ａ体積分率が大きく、圧縮降伏応力／引張降伏応力が低下している。鋼板Ｎｏ．１７は、Ｎｂの含有量が本発明の請求範囲よりも大きいため、析出強化にともなう靭性劣化のため、ＤＷＴＴ特性が低下している。鋼管Ｎｏ．１８は、本発明の必須添加元素であるＴｉが添加されていないため、加熱時の組織が粗大化し、それが圧延後の組織にも受け継がれるため、靱性が劣化し、ＤＷＴＴ特性が低下している。鋼板Ｎｏ．１９は、本発明の条件で加速冷却を実施した後、室温まで空冷後、炉加熱により再加熱を実施しているため、ベイナイトの焼戻しが過度であり、引張強さとＤＷＴＴ特性が劣化している。 Steel pipe No. In Nos. 15 and 16, since C, Mn, and Nb are larger than the claims of the present invention, the MA volume fraction is large, and the compressive yield stress / tensile yield stress is reduced. Steel plate No. In No. 17, since the Nb content is larger than the claimed range of the present invention, the DWTT characteristic is lowered due to toughness deterioration accompanying precipitation strengthening. Steel pipe No. No. 18, since Ti which is an essential additive element of the present invention is not added, the structure at the time of heating becomes coarse, and it is inherited by the structure after rolling, so that the toughness is deteriorated and the DWTT characteristic is lowered. Yes. Steel plate No. No. 19, after performing accelerated cooling under the conditions of the present invention, after air cooling to room temperature and then reheating by furnace heating, tempering of bainite is excessive, and tensile strength and DWTT characteristics are degraded. .
本発明により、圧潰強度および真円度に優れる石油や天然ガスの輸送とりわけ海底パイプラインに使用させる厚肉高強度ラインパイプ用として好適な厚鋼板を冷間で成形し溶接して製造される高靱性溶接鋼管の製造が可能となり、産業上極めて有効である。 According to the present invention, a high-temperature steel plate manufactured by cold forming and welding a thick steel plate suitable for transportation of oil and natural gas having excellent crushing strength and roundness, particularly for thick-walled high-strength line pipes used for submarine pipelines. This makes it possible to manufacture tough welded steel pipes and is extremely effective in the industry.
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| JP5055774B2 (en) * | 2005-03-17 | 2012-10-24 | Ｊｆｅスチール株式会社 | A steel plate for line pipe having high deformation performance and a method for producing the same. |
| US8715430B2 (en) * | 2005-04-04 | 2014-05-06 | Nippon Steel & Sumitomo Metal Corporation | High strength steel plate and high strength welded pipe excellent in ductile fracture characteristic and methods of production of same |
| JP4721956B2 (en) * | 2006-06-06 | 2011-07-13 | 株式会社神戸製鋼所 | Thick steel plate with excellent base metal toughness and fatigue crack growth characteristics |
| JP4940886B2 (en) * | 2006-10-19 | 2012-05-30 | Ｊｆｅスチール株式会社 | High strength steel plate for line pipe with excellent HIC resistance and method for producing the same |
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