JP2004076101A - High-strength high-toughness steel pipe material excellent in weldability and its production method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a steel pipe material which attains a high strength while satisfying a high absorption energy and is not hardened even when influenced by welding heat. <P>SOLUTION: The high-strength high-toughness steel pipe material is a steel plate comprising 0.005-0.020 mass% C, 0.05-1.0 mass% Si, 1.0-4.0 mass% Mn, 0.01-0.10 mass% Al, 0.01-0.50 mass% Nb, 0.005-0.10 mass% Ti, 0.0010-0.010 mass% B, 0.003 mass% or lower S, and the balance being Fe and unavoidable impurities, provided that X1 defined by equation (1): X1 = 970-130*Mn-1450*Nb is 400-650 and that the microstructure of the steel plate contains a bainite phase in the form of α°B in a content of at least 70 vol%. <P>COPYRIGHT: (C)2004,JPO

Description

【００４０】
【表２】

【００４１】
得られた鋼板からミクロ組織観察用の全厚×２０ｍｍ幅×１０ｍｍ高さのブロック試料をＬ断面（圧延方向に平行な板厚方向断面）が被検面となるように採取し、その被検面を３％ナイタール腐食液で処理してミクロ組織を現出させ、そのミクロ組織を走査型電子顕微鏡にて８００　〜２０００倍の適当な倍率で無作為に４視野以上写真撮影し、それぞれの写真中に観察されたα°_Ｂ形態のベイナイトの領域をトレース後、画像解析処理により前記トレース領域の全視野面積に対する面積率を計算し、ベイナイト組織が等方的形状であると仮定して（この仮定と実際との誤差は無視できる程度に小さいと考えられる。）、この計算値をα°_Ｂ形態のベイナイト相の体積率とした。この体積率を表２に示す。なお、フェライト、マルテンサイト、セメンタイトについても同様の方法で体積率を求めた。その値を表３に示す。なお、表３にはα°_Ｂ形態のベイナイト相の体積率も再掲した。
【００４２】
次に、上記の各鋼板から、ＪＩＳ　Ｚ　２２０１に規定されている４号引張試験片をＬ方向（圧延方向に平行な方向）が引張方向となるように採取し、ＪＩＳ　Ｚ　２２４１に規定されている引張試験を行い、０．２％耐力および引張強度を評価した。また、同鋼板からＪＩＳ　Ｚ　２２０２に規定されている４号シャルピー試験片をＣ方向（圧延幅方向に平行な方向）が試験片長手方向となるように採取し、ＪＩＳ　Ｚ　２２４２に規定されているシャルピー衝撃試験を行い、−４６　℃における吸収エネルギー（略号：ｖＥ_−４６）、および、脆性破面率の遷移曲線から５０％破面遷移温度（略号：ｖＴｒｓ）を評価した。
【００４３】
最後に、上記の各鋼板から、ＪＩＳ　Ｚ　３１５８に従ってｙ型溶接割れ試験（ｙスリット割れ試験）用試験体を採取・組立加工した後、環境温度１０℃、相対湿度８０％に設定した環境室内で１時間放置したＪＩＳ　Ｚ　３２１２に規定される低水素系溶接棒を乾燥処理せずに用いて、予熱温度１０℃とした試験体に試験ビードを溶接した。４８時間経過後、試験体の溶接部の５箇所から断面割れ観察用試料を切り出し、研削・研磨加工後に溶接部の割れを５倍の拡大投影機を用いて観察し、割れ長さを測定して断面割れ率を計算した。
【００４４】
上記の引張、衝撃試験結果およびｙスリット割れ試験結果を表３に示す。
【００４５】
【表３】

【００４６】
化学組成およびミクロ組織が本発明要件を満たし、該ミクロ組織が満たすべき本発明要件（：α°_Ｂ形態のベイナイト相≧７０ｖｏｌ．％　）が本発明に係る製造プロセスにより具現した発明例Ａ１〜Ｇ１、Ｓ１では、いずれも引張強度が７００Ｎ／ｍｍ^２を超える高強度でｖＥ_−４６も３００Ｊを超えるような高い吸収エネルギーを示した。また、ｙ型溶接割れ試験においても溶接部断面に割れは発生しなかった。
【００４７】
また、圧延終了温度が下限を下回った比較例Ｇ２および冷却開始温度が下限を下回った比較例Ｇ３は、いずれもミクロ組織観察でフェライト相が認められ、α°_Ｂの体積率が低下した結果、強度が低く、さらには同程度のｖＴｒｓであってもｖＥ_−４６も低い。特に、フェライト変態温度域で圧延していたＧ２では、セパレーションの発生で著しく吸収エネルギーが低下した。
【００４８】
また、冷却速度が下限を下回った比較例Ｇ４、および冷却停止温度が上限を上回った比較例Ｇ５は、ともにミクロ組織観察でα_Ｂ組織が多く、α°_Ｂの体積率が低下した結果、強度、ｖＥ_−４６ともＧ１に較べて低い。
一方、Ｘ２が６００　を上回った比較例Ｈ１も同様にα°_Ｂの体積率が低下し、強度、ｖＥ_−４６とも低い値となった。逆に、Ｘ２が４００　を下回った比較例Ｊ１は、強度は高いものの、ｙ型溶接割れ試験において、断面割れ率が８５％となり、実溶接施工では予熱が必須となる。これは、溶接部のミクロ組織がマルテンサイトとなり、非常に溶接割れを起こしやすくなったからである。
【００４９】
さらに、Ｃの上限を超えた比較例Ｋ１、Ｍｎの上限を超えた比較例Ｌ１も同様に、ｙ型溶接割れ試験で割れが発生した。また、比較例Ｋ１はα°_Ｂ中に第２相として島状マルテンサイトが多数観察され、この第２相の増加に伴いｖＥ_−４６が低下した。また、Ｎｂの上限を超えた比較例Ｍ１およびＴｉの上限を超えた比較例Ｑ１はいずれも析出に伴う硬化から脆性破壊が起こりやすくなり、ｖＴｒｓが上昇した結果、ｖＥ_−４６は低下した。また、Ｓの上限を超えた比較例Ｎ１とＢの下限を下回った比較例Ｒ１はいずれもα°_Ｂの体積率が７０％を下回っており、目標とする強度とｖＥ_−４６が得られなかった。
【００５０】
このように、本発明の鋼板では、従来なしえなかった高強度・高靭性と溶接時の低予熱温度でも割れが発生しないという優れた溶接性との両方を兼ね備えることができるようになった。
【００５１】
【発明の効果】
本発明によれば、炭素量の低減と、適切な合金元素添加と、適切な加熱‐熱間圧延‐加速冷却条件の組み合わせにより、α°_Ｂ形態のベイナイト体積率を７０％以上にすることにより、高強度かつ高シャルピー吸収エネルギーの鋼板特性と、溶接時の低予熱温度条件を許容する優れた溶接性とを具備する高強度鋼管素材が実現するという効果を奏する。[0001]
TECHNICAL FIELD OF THE INVENTION
TECHNICAL FIELD The present invention relates to a large-diameter welded steel pipe material used for pipelines or building structures, and particularly to a high-strength and toughness steel pipe material excellent in weldability, having a strength exceeding API-5LX80 class, and a method for producing the same.
[0002]
[Prior art]
In order to reduce oil pipeline laying costs, attempts have been made to reduce material costs by increasing the strength of steel pipes and reducing the pipe thickness. Conventionally, in a large-diameter welded steel pipe formed from a thick steel plate by a UOE process or a roll bender process, as described in JP-A-08-35011, elements such as Mn, Cu, Ni, Cr, Mo, and V are conventionally used. The steel with a large amount of steel is hot-rolled, and after the rolling, accelerated cooling is performed to increase the strength of the material thick steel plate. Also, in Japanese Patent Application Laid-Open No. 08-269544, ₁ ~ Ar ₃ Rolling is performed in a so-called two-phase region between temperatures, and after imparting work strengthening of ferrite, accelerated cooling is similarly performed to achieve higher strength.
[0003]
[Problems to be solved by the invention]
In recent years, studies on the safety evaluation of such high-strength steel pipes have been actively conducted, and at the same time, the ductile cracks are generated in the steel pipes by the action of sudden external force while preventing brittle fracture at the operating temperature. However, it has been required that the crack stop at a certain length to prevent the crack from propagating throughout the pipeline. It is known from research that the crack propagation arresting property is improved as the Charpy absorbed energy of the steel pipe base material is higher. In a high-strength steel pipe exceeding API-5LX80 class, a high absorption energy exceeding 300 J is obtained. Is estimated to be necessary.
[0004]
However, a technique for increasing the strength by a combination of an alloy element and accelerated cooling as disclosed in Japanese Patent Application Laid-Open No. 08-35011 cannot always stably achieve the Charpy high absorption energy of the base material. When work strengthening of ferrite is applied as disclosed in Japanese Patent Application Laid-Open No. 08-269544, separation (a fracture surface becomes almost parallel to a rolling surface) is caused in a specimen during a Charpy test due to a texture formed in the ferrite. Absorbed energy is rather reduced due to the occurrence of brittle fracture). As described above, means for achieving high strength while satisfying high absorption energy has not been clarified.
[0005]
In addition, in the construction of pipelines, pipe-to-pipe joining means is generally covered gas arc welding (SMAW) on site, but this welding method tends to involve moisture from the environment into the weld metal, and as a result, the weld metal Hydrogen diffusing from the base material toward the base metal gathers in a hardened region under the influence of welding and is liable to cause so-called delayed cracking. For this reason, so-called pre-heat treatment is required to promote the diffusion of hydrogen and allow the hydrogen to escape to the outside of the steel sheet. Such pre-heat treatment increases the welding time and deteriorates the working environment of the welding worker. Therefore, it is desired that the construction can be performed without preheating, even for a high-strength pipe. That is, as a countermeasure on the steel sheet side, it is desired that the steel sheet does not harden even under the influence of welding heat.
[0006]
The present invention, in view of the current state of the prior art, achieves high strength while satisfying high absorption energy, does not harden even under the influence of welding heat, and provides a high strength and high toughness steel pipe material excellent in weldability and a method of manufacturing the same. The purpose is to provide.
[0007]
[Means for Solving the Problems]
The present inventors have conducted intensive studies on increasing the strength by controlling the microstructure, and making the microstructure of the material steel sheet bainite, compared with the case where the structure control such as ferrite-bainite or ferrite-martensite was performed, And a good balance of Charpy absorbed energy. Furthermore, focusing on the second phase of the bainite structure, and by reducing these second phase structures, the tensile strength of bainite is reduced, but a high Charpy absorbed energy exceeding 300 J at -46 ° C is achieved. I found it. The reduction of the second phase in the bainite is performed by setting the carbon content of the steel to 0.02 mass% or less, which is the solid solubility limit of bcc iron, so that the diffusion transfer and enrichment of C do not occur during the transformation from austenite to bainite. Can be achieved.
[0008]
Next, as a result of continuing research on the establishment of a technique for increasing the strength of the bainite structure in which the second phase structure has been reduced as much as possible, it was found that a combination of hardenability improving elements such as Mn, Cu, Ni, Cr, Mo, and Nb, By combination with the subsequent cooling, bainitic ferrite (symbol α °) _B It was found that the strength increased by increasing the volume fraction of the bainite structure in the form called ")." Furthermore, this α ° _B The phase fraction of the bainite structure exhibiting a morphology is 70 vol. It was also found that when the amount is increased to more than%, the austenite before transformation is strongly worked in a low temperature range to inherit the strain introduced into the austenite, so that the strength can also be increased depending on the controlled rolling conditions during hot rolling. Α ° by strong austenitic processing _B The high strength of the bainite in the form maintains the Charpy high absorption energy characteristic of the bainite excluding the second phase structure, except when the austenite processing temperature range becomes too low and separation occurs in the transformed bainite. I do.
[0009]
By the above alloy element adjustment, hot rolling and accelerated cooling control after hot rolling, the problem of high strength and high Charpy absorbed energy was solved.
Next, when various evaluations were made on the weld cracking resistance during welding when the strength was increased by such means, α ° without martensitic transformation was observed in the heat history during welding. _B It has been found that when bainite transformation occurs, the hardness of the weld heat affected zone is suppressed from increasing, and delayed cracking does not occur even under a low preheating condition of 10 ° C. Specifically, in the combination of the added amounts of Mn, Cu, Ni, Cr, Mo, and Nb (the content by addition, the same applies hereinafter), the transformed structure after hot rolling and subsequent accelerated cooling is α ° _B The bainite phase in the form of 70 vol. It is thought that it is sufficient to determine a range that does not result in martensite structure in the heat history during welding, and regression analysis of the experimental results shows that this range is defined by the following equation (2). Can be well described by the index X2,
X2 = 970-130 * Mn-55 * Cu-30 * Ni-70 * Cr-90 * Mo-1450 * Nb (2)
In order to suppress martensitic transformation, X2 ≧ 400 and α ° _B The bainite transformation in the form of 70 vol. %, It was found that X2 ≦ 650 should be satisfied.
[0010]
The present invention has been made based on the above findings, and the gist is as follows.
(1) C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0 0.10 to 0.50 mass%, Ti: 0.005 to 0.10 mass%, B: 0.0010 to 0.010 mass%, S: 0.003 mass% or less, and the following formula (1) X1 defined is 400 to 650, and a steel sheet as a steel pipe material comprising the balance of Fe and unavoidable impurities, wherein the microstructure of the steel sheet is α ° _B The bainite phase in the form of 70 vol. % High-strength, high-toughness steel pipe material with excellent weldability characterized by containing at least 0.1%.
[0011]
X1 = 970-130 * Mn-1450 * Nb (1)
(2) C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0 0.01 to 0.50 mass%, Ti: 0.005 to 0.10 mass%, B: 0.0010 to 0.010 mass%, S: 0.003 mass% or less, and Cu: 0.5 to 0.5 mass%. 3.0 mass%, Ni: 0.2 to 3.0 mass%, Cr: 0.2 to 1.0 mass%, Mo: 0.1 to 1.0 mass%, or two or more, and / or , Ca: 0.001 to 0.020 mass%, REM: one or two of 0.005 to 0.020 mass%, and X2 defined by the following formula (2) is 400 to 650. And the balance Fe And a steel sheet as a steel pipe material comprising unavoidable impurities, wherein the microstructure of the steel sheet is α ° _B The bainite phase in the form of 70 vol. % High-strength, high-toughness steel pipe material with excellent weldability characterized by containing at least 0.1%.
[0012]
X2 = 970-130 * Mn-55 * Cu-30 * Ni-70 * Cr-90 * Mo-1450 * Nb (2)
(3) C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0 0.01 to 0.50 mass%, Ti: 0.005 to 0.10 mass%, B: 0.0010 to 0.010 mass%, S: 0.003 mass% or less,
Alternatively, one of Cu: 0.5 to 3.0 mass%, Ni: 0.2 to 3.0 mass%, Cr: 0.2 to 1.0 mass%, and Mo: 0.1 to 1.0 mass%. And / or two or more of Ca and 0.001 to 0.020 mass%, and REM: 0.005 to 0.020 mass%.
In addition, X2 defined by the following formula (2) is 400 to 650, and a steel slab consisting of the balance Fe and unavoidable impurities is heated to 1000 to 1250 ° C. and then hot-rolled into a steel sheet. The cumulative rolling reduction in the low-temperature austenite temperature range of 900 ° C or less is set to 50% or more, the rolling end temperature is set to 700 to 850 ° C, and then the steel sheet is cooled from the temperature of -50 ° C or more to the cooling rate of 5 ° C / s. A method for producing a high-strength, high-toughness steel pipe material excellent in weldability, characterized by water-cooling to a temperature of 400 ° C. or less as described above.
[0013]
X2 = 970-130 * Mn-55 * Cu-30 * Ni-70 * Cr-90 * Mo-1450 * Nb (2)
In the formulas (1) and (2), each element symbol indicates the content (mass%) of the element in steel, “*” indicates a product operator, and “−” indicates a difference operator.
[0014]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the reasons for limiting the chemical composition (chemical component content), microstructure, and manufacturing process (heating, hot rolling, accelerated cooling) in the present invention as described above will be described.
First, the reasons for limiting the chemical composition will be described.
[0015]
C: 0.005 to 0.020 mass%
The C content affects the formation of the second phase that lowers the Charpy absorbed energy of the steel sheet having the bainite structure. By setting the C content to 0.020 mass% or less, the formation of the second phase can be substantially suppressed, and a high absorption energy exceeding 300 J can be achieved. Therefore, the upper limit is set to 0.020 mass%. On the other hand, even if the extremely low carbon content of less than 0.005 mass% is performed, further improvement in Charpy absorbed energy is not expected and the cost during steelmaking only increases, so the lower limit was made 0.005 mass%.
[0016]
Si: 0.05 to 1.0 mass%
Si is required to be 0.05 mass% or more in steel production, and increases the strength of the steel by solid solution strengthening with an increase in the addition amount. However, if added in excess of 1.0 mass%, the base material is liable to cause brittle fracture at low temperatures, so the upper limit was set to 1.0 mass%. The preferred range is 0.1 to 0.5 mass%.
[0017]
Mn: 1.0 to 4.0 mass%
Mn is an element that enhances the hardenability, and when added according to formula (1) or formula (2) described below, the form of bainite changes to α ° _B It can be. Further, since it is inexpensive as compared with others, by setting the lower limit to 1.0 mass%, it is possible to suppress the increase in cost and increase the strength. However, if added in excess of 4.0 mass%, martensitic transformation of the weld is caused to promote delayed cracking of the weld, so the upper limit was made 4.0 mass%. The preferred range is 1.5 to 2.5 mass%.
[0018]
Al: 0.01 to 0.10 mass%
Al is added as a deoxidizing agent at the time of steel making. However, a small amount of Al added to a steel sheet is less than 0.01 mass%, which tends to cause insufficient deoxidation. Therefore, the lower limit is set to 0.01 mass%. On the other hand, if added in excess of 4.0 mass%, the cleanliness of the base material deteriorates and the energy absorbed by Charpy decreases, so the upper limit was set to 0.10 mass%.
[0019]
Nb: 0.01 to 0.50 mass%
Nb is required to be at least 0.01 mass% in order to extend the austenite non-recrystallization temperature range to a higher temperature side. Further, by adding according to the formula (1) or the formula (2) described later, the form of bainite is changed to α ° _B It can be. Due to the effect of this Nb addition (the inheritance of processing strain introduced by rolling at 900 ° C. or less), α ° _B The bainite having the form is further strengthened. On the other hand, if added in excess of 0.50 mass%, the base material is liable to cause brittle fracture at low temperatures, so the upper limit was set to 0.50 mass%. The preferred range is 0.015 to 0.06 mass%.
[0020]
Ti: 0.005 to 0.10 mass%
Ti is required to be 0.005% by mass or more in order to fix inevitably present free N in steel as TiN. This TiN also contributes to suppressing austenite grain growth in the heat affected zone. On the other hand, if added in excess of 0.10 mass%, the excess Ti forms carbides, significantly increasing the strength of the steel and easily causing brittle fracture. Therefore, the upper limit was set to 0.10 mass%. The preferred range is 0.005 to 0.020 mass%.
[0021]
B: 0.0010 to 0.010 mass%
B has the effect of suppressing the ferrite transformation from the austenite grain boundaries and facilitating the bainite transformation during the transformation that occurs during the cooling process after hot rolling. In particular, in the present invention, since the amount of C is reduced, 0.0010 mass% or more is required to suppress ferrite transformation. On the other hand, even if added in excess of 0.010 mass%, the effect is saturated, so the upper limit was set to 0.010 mass%. The preferred range is 0.0015 to 0.0030 mass%.
[0022]
S: 0.003 mass% or less
S is inevitably mixed into steel as an impurity element, but exists in the steel as MnS unless the form control or the like is performed. Since MnS is likely to become a transformation nucleus of ferrite and causes ferrite to be formed prior to bainite transformation, it is necessary to reduce the amount of S to reduce the amount of MnS. Therefore, the upper limit of the amount of S is set to 0.003 mass%. . When the form control by adding Ca or REM is not performed, it is preferable to reduce the amount to less than 0.0010 mass%.
[0023]
X1: 400 to 650
In the present invention, when Cu, Ni, Cr, and Mo are not added, X1 defined by the formula (1) (the formula (2) with the terms Cu, Ni, Cr, and Mo removed) is 400. The amounts of Mn and Nb are adjusted so that ≦ X1 ≦ 650. By setting X1 ≧ 400, martensitic transformation under the heat history during welding can be suppressed, and the occurrence of delayed cracking can be suppressed even at the lower limit preheating temperature of 10 ° C. On the other hand, by setting X1 ≦ 650, α ° can be obtained within the practical control range of the hot rolling and cooling conditions. _B The bainite phase in the form of 70 vol. % Or more, a steel sheet having a microstructure containing at least% can be obtained, and high strength of the steel sheet can be achieved.
[0024]
X2: 400 to 650
In the present invention, one or more selected from Cu, Ni, Cr, and Mo can be added as necessary, in addition to the component elements limited as described above. In this case, the amounts of Mn and Nb, and further, the amounts of Cu, Ni, Cr and Mo are adjusted so that X2 defined by the formula (2) satisfies 400 ≦ X2 ≦ 650. By setting X2 ≧ 400, martensitic transformation under the heat history during welding can be suppressed, and the occurrence of delayed cracking can be suppressed even at the lower limit preheating temperature of 10 ° C. On the other hand, by setting X2 ≦ 650, α ° can be obtained within the practical control range of the hot rolling and cooling conditions. _B The bainite phase in the form of 70 vol. % Or more, a steel sheet having a microstructure containing at least% can be obtained, and high strength of the steel sheet can be achieved.
[0025]
However, when Cu, Ni, Cr, and Mo are added, the content of each component is preferably in the following range.
Cu: 0.5 to 3.0 mass%
Cu is added at 0.5 mass% or more to α ° _B Although it contributes to morphology, if it is added in excess of 3.0 mass%, brittle fracture of the base material is likely to occur due to the precipitate dispersing effect. Therefore, the upper limit is set to 3.0 mass%. The preferred range is 0.05 to 1.50 mass%.
[0026]
Ni: 0.2 to 3.0 mass%
Ni is added at 0.2 mass% or more to α ° _B Contributes to the promotion of On the other hand, even if added in excess of 3.0 mass%, the effect is saturated, so the upper limit was set to 3.0 mass%. In addition, a suitable range is 0.25 to 1.0 mass%.
Cr: 0.2 to 1.0 mass%
Cr is added at 0.2 mass% or more to α ° _B Contributes to the promotion of On the other hand, if added in excess of 1.0 mass%, brittle fracture of the base material is likely to occur, so the upper limit was set to 1.0 mass%. The preferred range is 0.25 to 0.60 mass%.
[0027]
Mo: 0.1 to 1.0 mass%
Mo is α ° with addition of 0.1 mass% or more. _B Contributes to the promotion of On the other hand, if added in excess of 1.0 mass%, the precipitation dispersion strengthening of Mo carbides becomes excessive and brittle fracture is likely to occur, so the upper limit was set to 1.0 mass%. The preferred range is 0.1 to 0.6 mass%. In the present invention, for the purpose of controlling the form of inclusions, one or two selected from Ca and REM can be added in the following component content ranges.
[0028]
Ca: 0.001 to 0.020 mass%
If Ca is inevitably present in the steel and non-metallic inclusions MnS pose a problem in HAZ toughness, etc., by adding 0.001 mass% or more, the form of inclusions in CaS 2 generated at higher temperature is controlled. , The effect can be eliminated. However, if it is added in excess of 0.020 mass%, CaS 2 is formed in a cluster and has a rather bad effect, so the upper limit was made 0.020 mass%.
[0029]
REM: 0.005 to 0.020 mass%
In the case where non-metallic inclusions MnS unavoidably present in steel poses a problem in HAZ toughness, REM controls the inclusion morphology in REM sulfides generated at higher temperatures by adding 0.005 mass% or more. To eliminate that effect. However, if added in excess of 0.020 mass%, the cleanliness of the steel is degraded, so the upper limit was made 0.020 mass%.
[0030]
Next, the reasons for limiting the microstructure of the steel sheet will be described.
α ° _B (Bainitic ferrite) form of bainite phase ≧ 70 vol. %
The bainite structure of steel with a low carbon content has a morphology of α _B (Granular bainiticferrite) and α ° _B (Α _B , Α ° _B For details on the form, see "Bainite Research Group of the Iron and Steel Institute of Japan, Bainite Research and Editing Group: Bainite Photographs of Steel-1, ---- Continuous Cooling (Intermediate Stage) Transformation Structure of Low Carbon Steel ---, June 1992 , Page 24 "). Α ° _B The bainite structure having a morphology has a fraction of 70 vol. % Or more, the austenite before transformation is subjected to strong working in a low temperature range to inherit the strain introduced into the austenite, so that the strength can be increased depending on the controlled rolling conditions during hot rolling, and such austenite can be increased. Since high Charpy absorbed energy exceeding 300 J at -46 ° C can be achieved even at high strength by strong working of _B The bainite structure of the form is 70 vol. % Or more. The martensite or cementite produced as a phase other than the bainite structure of the steel obtained in the present invention is 2 vol. % Or less, and it is considered that the fact that the amount of martensite or cementite is small leads to an improvement in Charpy absorbed energy.
[0031]
Next, the manufacturing process will be described.
In the manufacturing process according to the present invention, a steel slab having a composition within the above-mentioned limited range is formed into a product steel plate by a manufacturing process including a sequential process of heating, hot rolling, and accelerated cooling. Shall be satisfied.
Heating temperature: 1000-1250 ° C
When the heating temperature of the slab is 1000 ° C. or higher, uniform austenite is obtained. Therefore, the lower limit of the heating temperature is 1000 ° C. On the other hand, when heated to more than 1250 ° C., the austenite grains became extremely coarse, and when hot rolling was performed as it was, the toughness of the steel sheet was significantly reduced. In addition, more preferably, it is 1050-1150 degreeC.
[0032]
Cumulative rolling reduction in low temperature austenite region below 900 ° C ≧ 50%
The heated slab is immediately subjected to hot rolling. In particular, the rolling reduction in the so-called austenite unrecrystallized region at 900 ° C or lower is performed by a rolling schedule such that the cumulative rolling reduction becomes 50% or more, thereby increasing the cumulative rolling reduction. With α ° _B The strength of the bainite in the form is increased, and a desired high strength can be achieved. Therefore, the cumulative rolling reduction at 900 ° C. or less in hot rolling is set to 50% or more.
[0033]
Hot rolling end temperature: 700 to 850 ° C
In rolling at a low temperature where austenite does not recrystallize, the lower the rolling temperature, the greater the strain accumulation effect. However, when the temperature is kept below 700 ° C., a rolled texture is formed in austenite, which causes Has a strong tendency to cause separation, and the Charpy absorbed energy is significantly reduced. Therefore, the lower limit of the rolling end temperature was set to 700 ° C. On the other hand, if the rolling end temperature is higher than 850 ° C., it becomes difficult to secure the above-mentioned cumulative rolling reduction of 50% or less at 900 ° C. or less in actual operation. Therefore, the upper limit of the rolling end temperature was set to 850 ° C.
[0034]
Cooling (water cooling) start temperature ≧ rolling end temperature−50 ° C.
The hot-rolled product (steel plate) needs to be water-cooled as soon as possible after rolling is completed (by minimizing the time of air-cooling until water-cooling) in order to transform it into bainite. In particular, when water cooling is started after the temperature of the steel sheet falls below the rolling end temperature of −50 ° C., ferrite transformation occurs between the end of rolling and the start of water cooling, leading to a decrease in YS and TS due to ferrite generation. The rolling end temperature was -50 ° C or higher.
[0035]
Cooling rate ≧ 5 ℃ / s
By optimizing the amounts of Mn, Cu, Ni, Cr, Mo, and Nb, if a cooling rate of 5 ° C./s or more is secured in water cooling after hot rolling within the above-described limited range according to the manufacturing process conditions, Α ° to aim bainite transformation without ferrite transformation _B Since a bainite structure in a form is obtained, the lower limit of the cooling rate in water cooling after hot rolling was set at 5 ° C./s. Although the upper limit of the cooling rate is not particularly set, the maximum cooling rate that can be practically used is 50 ° C./s, and is preferably 5 to 50 ° C./s.
[0036]
Cooling stop temperature ≤ 400 ° C
In the alloy element design in the present invention, the transformation end temperature in the continuous cooling transformation is considered to be 400 ° C. or higher. Therefore, it is 400 ° C. at the minimum that austenite completely forms a bainite structure, and it is sufficient to continue water cooling to a temperature not exceeding 400 ° C., so the upper limit of the cooling stop temperature is set to 400 ° C.
[0037]
Incidentally, the slab to be subjected to the production process according to the present invention, the production method is not particularly limited, according to a conventional method, after melting the steel by the open-hearth method, the converter method or the electric furnace method, component adjustment, The casting may be performed by any of the continuous casting method and the ingot making method. Further, in forming the manufactured steel sheet into a steel pipe, high strength and high absorption energy, and high resistance to delayed welding cracking, which are the objects of the present invention, are achieved regardless of whether the UOE process or the roll bender process is used. .
[0038]
【Example】
Using steel slabs having the chemical composition shown in Table 1, thick steel plates having a thickness of 15 to 30 mm were manufactured under the heating-hot rolling-cooling conditions shown in Table 2.
[0039]
[Table 1]

[0040]
[Table 2]

[0041]
From the obtained steel sheet, a block sample of total thickness × 20 mm width × 10 mm height for microstructure observation is sampled so that an L section (a section in a thickness direction parallel to the rolling direction) is a test surface, and the test is performed. The surface was treated with a 3% nital etching solution to reveal a microstructure, and the microstructure was photographed at random at an appropriate magnification of 800 to 2000 times with a scanning electron microscope in four or more visual fields. Α ° observed during _B After tracing the bainite region of the form, the area ratio of the trace region with respect to the entire visual field area is calculated by image analysis processing, and it is assumed that the bainite structure has an isotropic shape (an error between this assumption and the actual one is ignored. It is considered to be as small as possible.) _B The volume fraction of the bainite phase was determined. Table 2 shows this volume ratio. The volume ratio of ferrite, martensite and cementite was determined in the same manner. Table 3 shows the values. Table 3 shows α ° _B The volume fraction of the bainite phase in the morphology is also given again.
[0042]
Next, a No. 4 tensile test piece specified in JIS Z 2201 was sampled from each of the above steel sheets so that the L direction (direction parallel to the rolling direction) was the tensile direction, and the tensile test piece was specified in JIS Z 2241. A tensile test was performed to evaluate 0.2% proof stress and tensile strength. A No. 4 Charpy test piece specified in JIS Z 2202 is sampled from the steel sheet so that the C direction (direction parallel to the rolling width direction) is the longitudinal direction of the test piece, and is specified in JIS Z 2242. A Charpy impact test was performed, and the absorbed energy at −46 ° C. (abbreviation: vE ₋₄₆ ) And a 50% fracture transition temperature (abbreviation: vTrs) was evaluated from the transition curve of the brittle fracture ratio.
[0043]
Lastly, a specimen for y-type welding cracking test (y-slit cracking test) is sampled and assembled from each of the above-mentioned steel sheets in accordance with JIS Z 3158, and then is placed in an environmental chamber set at an environmental temperature of 10 ° C. and a relative humidity of 80%. A test bead was welded to a test specimen at a preheating temperature of 10 ° C. using a low hydrogen welding rod specified in JIS Z 3212 which had been left for 1 hour without drying treatment. After a lapse of 48 hours, samples for observing cross-sectional cracks were cut out from five places in the welded portion of the test specimen, and after grinding and polishing, the cracks in the welded portion were observed using a magnifying projector of 5 times, and the crack length was measured. Then, the section cracking rate was calculated.
[0044]
Table 3 shows the results of the tensile and impact tests and the y-slit crack test.
[0045]
[Table 3]

[0046]
The chemical composition and the microstructure satisfy the requirements of the present invention, and the requirements of the present invention to be satisfied by the microstructure (: α ° _B Form bainite phase ≧ 70 vol. %), In the invention examples A1 to G1 and S1 embodied by the manufacturing process according to the invention, the tensile strength is 700 N / mm. ² VE with high strength exceeding ₋₄₆ Also showed a high absorption energy exceeding 300 J. Also, no crack was generated in the welded section in the y-type weld crack test.
[0047]
In Comparative Example G2 in which the rolling end temperature was lower than the lower limit and in Comparative Example G3 in which the cooling start temperature was lower than the lower limit, a ferrite phase was observed in microstructure observation, and α ° _B As a result, the strength is low and vE ₋₄₆ Is also low. In particular, in G2 which was rolled in the ferrite transformation temperature range, the absorption energy was significantly reduced due to the occurrence of separation.
[0048]
In Comparative Example G4 in which the cooling rate was lower than the lower limit, and in Comparative Example G5 in which the cooling stop temperature was higher than the upper limit, both of the microstructures were observed by α. _B Many tissues, α ° _B Of the strength, vE ₋₄₆ Both are lower than G1.
On the other hand, Comparative Example H1 in which X2 exceeded 600 1 _B Volume ratio decreases, strength, vE ₋₄₆ Both were low values. Conversely, Comparative Example J1 in which X2 was less than 400 had high strength, but had a cross-sectional cracking rate of 85% in the y-type welding crack test, and required preheating in actual welding work. This is because the microstructure of the welded portion became martensite, and welding cracks were extremely likely to occur.
[0049]
Furthermore, the comparative example K1 which exceeded the upper limit of C and the comparative example L1 which exceeded the upper limit of Mn also similarly cracked in the y-type welding crack test. Comparative Example K1 is α ° _B A large number of island-like martensite was observed as the second phase, and with the increase of this second phase, vE ₋₄₆ Decreased. Further, in each of Comparative Example M1 exceeding the upper limit of Nb and Comparative Example Q1 exceeding the upper limit of Ti, brittle fracture was liable to occur due to hardening caused by precipitation, and vTrs increased. ₋₄₆ Fell. Further, Comparative Example N1 which exceeded the upper limit of S and Comparative Example R1 which fell below the lower limit of B were all α °. _B Is less than 70%, the target strength and vE ₋₄₆ Was not obtained.
[0050]
As described above, the steel sheet of the present invention can have both high strength and high toughness, which could not be achieved conventionally, and excellent weldability in which cracks do not occur even at a low preheating temperature during welding.
[0051]
【The invention's effect】
According to the present invention, by reducing the amount of carbon, adding an appropriate alloy element, and combining appropriate heating-hot rolling-accelerated cooling conditions, α ° _B By making the bainite volume fraction of the form 70% or more, a high strength steel pipe material with high strength and high Charpy absorbed energy steel sheet properties and excellent weldability that allows low preheating temperature conditions during welding is realized. It has the effect of doing.

Claims (3)

C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0.01 to 0.50 mass%, Ti: 0.005 to 0.10 mass%, B: 0.0010 to 0.010 mass%, S: 0.003 mass% or less, and further defined by the following formula (1) X1 is 400-650, a steel plate as a steel pipe material and the balance Fe and unavoidable impurities, 70 vol bainite phase in the form of the microstructure of the steel plate is alpha ° _B. % High-strength, high-toughness steel pipe material with excellent weldability characterized by containing at least 0.1%.
X1 = 970-130 * Mn-1450 * Nb (1) C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0.01 to 0.50% by mass, Ti: 0.005 to 0.10% by mass, B: 0.0010 to 0.010% by mass, S: 0.003% by mass or less, and further, Cu: 0.5 to 3.0% by mass. %, Ni: 0.2 to 3.0 mass%, Cr: 0.2 to 1.0 mass%, Mo: 0.1 to 1.0 mass%, and / or Ca: 0.001 to 0.020 mass%, REM: contains one or two of 0.005 to 0.020 mass%, and X2 defined by the following formula (2) is 400 to 650; Balance Fe and A steel sheet as the steel pipe material consisting of unavoidable impurities, 70 vol bainite phase in the form of the microstructure of the steel plate is alpha ° _B. % High-strength, high-toughness steel pipe material with excellent weldability characterized by containing at least 0.1%.
X2 = 970-130 * Mn-55 * Cu-30 * Ni-70 * Cr-90 * Mo-1450 * Nb (2) C: 0.005 to 0.020 mass%, Si: 0.05 to 1.0 mass%, Mn: 1.0 to 4.0 mass%, Al: 0.01 to 0.10 mass%, Nb: 0.01 to 0.50% by mass, Ti: 0.005 to 0.10% by mass, B: 0.0010 to 0.010% by mass, S: 0.003% by mass or less,
Alternatively, one of Cu: 0.5 to 3.0 mass%, Ni: 0.2 to 3.0 mass%, Cr: 0.2 to 1.0 mass%, and Mo: 0.1 to 1.0 mass%. And / or two or more of Ca and 0.001 to 0.020 mass%, and REM: 0.005 to 0.020 mass%.
In addition, X2 defined by the following formula (2) is 400 to 650, and a steel slab consisting of the balance Fe and unavoidable impurities is heated to 1000 to 1250 ° C. and then hot-rolled into a steel sheet. The cumulative rolling reduction in the low-temperature austenite temperature range of 900 ° C or less is set to 50% or more, the rolling end temperature is set to 700 to 850 ° C, and then the steel sheet is cooled from the temperature of -50 ° C or more to the cooling rate of 5 ° C / s. A method for producing a high-strength, high-toughness steel pipe material excellent in weldability, characterized by water-cooling to a temperature of 400 ° C. or less as described above.
X2 = 970-130 * Mn-55 * Cu-30 * Ni-70 * Cr-90 * Mo-1450 * Nb (2)