JP2005146407A - Ultrahigh strength steel sheet and ultrahigh strength steel tube having excellent high speed ductile fracture property, and their production method - Google Patents
Ultrahigh strength steel sheet and ultrahigh strength steel tube having excellent high speed ductile fracture property, and their production method Download PDFInfo
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- 229910000797 Ultra-high-strength steel Inorganic materials 0.000 title claims abstract description 41
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- 229910000831 Steel Inorganic materials 0.000 claims abstract description 63
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Abstract
Description
本発明は、天然ガス・原油輸送用ラインパイプ等に好適な、800MPa以上の引張強さ(TS)を有する高速延性破壊特性に優れた超高強度ラインパイプに関する。 The present invention relates to an ultra-high-strength line pipe excellent in high-speed ductile fracture characteristics having a tensile strength (TS) of 800 MPa or more, which is suitable for a natural gas / crude oil transportation line pipe or the like.
近年、原油・天然ガスのパイプラインにおいて、輸送効率の向上を目的とした高内圧化や現地施工能率の向上を目的としたラインパイプの外径、重量の低減が要求され、X100(引張強さ760MPa以上)を超える高強度鋼管の開発が進められている(例えば、特許文献1、2)。
In recent years, in pipelines for crude oil and natural gas, it has been required to reduce the outer diameter and weight of the line pipe for the purpose of increasing the internal pressure for the purpose of improving transportation efficiency and improving the efficiency of local construction. Development of high-strength steel pipes exceeding 760 MPa (for example,
また、パイプラインでは、鋼管の母材に発生した延性亀裂が管軸方向に100m/s以上もの高速で、100mから数kmにも及ぶ長距離を伝播する可能性があり、耐アレスト性が要求される。耐アレスト性は、亀裂の伝播を停止させる特性であり、脆性亀裂が母材を伝播して停止する特性、即ち耐脆性破壊特性と、延性亀裂が母材を伝播して停止する特性、即ち高速延性破壊特性に分類される。 Also, in pipelines, ductile cracks that occur in the base material of steel pipes can propagate over a long distance ranging from 100 m to several km at a high speed of 100 m / s or more in the pipe axis direction, requiring arrest resistance. Is done. The arrest resistance is a property that stops the propagation of cracks, a property that a brittle crack propagates and stops the base material, that is, a brittle fracture resistance property, and a property that a ductile crack propagates and stops the base material, that is, a high speed. Classified as ductile fracture characteristics.
耐脆性破壊特性は、落重破壊試験(Drop Weight Tear Test、DWTT試験という)を行い、延性破面率が85%以上になる温度(DWTT遷移温度という)で評価される。脆性亀裂は溶接部から発生することが多く、試験片の中央部に溶接ビードを形成して脆性亀裂を導入し、DWTT試験を行って評価した耐脆性破壊特性に優れた鋼管が提案されている(例えば、特許文献3)。 Brittle fracture characteristics, drop weight destructive test (D rop W eight T ear T est, referred DWTT test) performed, ductile fracture rate is evaluated at a temperature equal to or higher than 85% (referred DWTT transition temperature). Brittle cracks often occur from welds, and steel pipes with excellent brittle fracture resistance evaluated by performing a DWTT test by introducing a weld bead at the center of the specimen and introducing the brittle cracks have been proposed. (For example, patent document 3).
これに対して、高速延性破壊特性の評価には、鋼管の表面に爆薬を装着後、爆発させて発生した延性亀裂が停止するか否かを判定するフルクラックバーストテストが最適である。しかし、フルクラックバーストテストは、試験に要するコストが非常に高いため、従来、フルクラックバーストテストの結果と比較的よく一致するシャルピー吸収エネルギー又はDWTT試験によって求められる吸収エネルギー(DWTT吸収エネルギーという)で評価されていた。 On the other hand, the full crack burst test for determining whether or not the ductile cracks generated by explosion after the explosive is mounted on the surface of the steel pipe is optimal for the evaluation of the high-speed ductile fracture characteristics. However, since the cost required for the full crack burst test is very high, conventionally, the Charpy absorbed energy or the absorbed energy required by the DWTT test (referred to as DWTT absorbed energy) that is relatively well matched with the result of the full crack burst test. It was evaluated.
しかし、フルクラックバーストテストとシャルピー吸収エネルギー及びDWTT吸収エネルギーに良い相関が認められていたのは、X65程度の強度を有する鋼材であり、X100超の超高強度鋼では、高速延性破壊特性の評価にシャルピー衝撃試験及びDWTT試験が適していないことがわかった。そのため、試験コストが高いフルクラックバーストテストの代替として、高速延性破壊特性を簡便に評価し得る試験方法が必要とされ、その試験によって得られた知見を活用し、高速延性破壊特性に優れた鋼管の開発が要望されていた。 However, a good correlation was observed between the full crack burst test and Charpy absorbed energy and DWTT absorbed energy, which is a steel material having a strength of about X65. For ultra-high strength steel exceeding X100, evaluation of high-speed ductile fracture characteristics The Charpy impact test and the DWTT test were not suitable. Therefore, a test method that can easily evaluate high-speed ductile fracture characteristics is required as an alternative to the full-crack burst test, which has a high test cost. Steel pipes that have excellent high-speed ductile fracture characteristics by utilizing the knowledge gained from these tests The development of was requested.
本発明は、高速延性破壊特性を適正に評価し得る簡便な試験方法を確立し、それにより得た知見に基づいて、高速延性破壊特性に優れた引張強さ800MPa以上(API規格X100以上)の超高強度鋼管及びその製造方法を提供するものである。 The present invention establishes a simple test method capable of appropriately evaluating high-speed ductile fracture characteristics, and based on the knowledge obtained thereby, has a tensile strength of 800 MPa or more (API standard X100 or more) excellent in high-speed ductile fracture characteristics. An ultra-high strength steel pipe and a method for producing the same are provided.
本発明者は、引張強さが800MPa以上の超高強度鋼管の高速延性破壊特性を適正に評価し得る簡便な試験方法について検討を行い、更に高速延性破壊特性に優れた超高強度鋼管を得るための母材の成分及びミクロ組織について検討を行い、ミクロ組織、集合組織を最適化することが有効であるという知見を得、更に製造条件について検討を行い、高速延性超高強度鋼板及び鋼管並びにそれらの製造方法を発明するに至った。本発明の要旨は以下のとおりである。 The present inventor examines a simple test method capable of appropriately evaluating the high-speed ductile fracture characteristics of an ultra-high-strength steel pipe having a tensile strength of 800 MPa or more, and further obtains an ultra-high-strength steel pipe excellent in high-speed ductile fracture characteristics. To study the components and microstructure of the base material for the purpose, to obtain the knowledge that it is effective to optimize the microstructure and texture, to further study the production conditions, high-speed ductile ultra-high strength steel sheet and steel pipe, It came to invent those manufacturing methods. The gist of the present invention is as follows.
(1)ミクロ組織が、面積率で1〜5%のフェライト及び残部がベイナイト・マルテンサイトからなり、圧延方向を軸として圧延面から45°回転させた断面の(100)の集積度が3以下であることを特徴とする高速延性破壊特性に優れた超高強度鋼板。 (1) The microstructure has an area ratio of 1 to 5% of ferrite and the balance is bainite martensite, and the degree of integration of (100) in the section rotated 45 ° from the rolling surface with the rolling direction as the axis is 3 or less. An ultra-high-strength steel sheet with excellent high-speed ductile fracture characteristics.
(2)ミクロ組織が、面積率で5%超〜60%のフェライト及び残部がベイナイト・マルテンサイトからなり、圧延方向を軸として圧延面から45°回転させた断面の(100)の集積度が3以下であることを特徴とする高速延性破壊特性に優れた超高強度鋼板。 (2) The microstructure has an area ratio of more than 5% to 60% of ferrite and the balance of bainite / martensite, and the degree of integration of (100) in the cross section rotated 45 ° from the rolling surface with the rolling direction as the axis. An ultra-high-strength steel sheet excellent in high-speed ductile fracture characteristics, characterized by being 3 or less.
(3)フェライトの平均粒径が5μm以下であることを特徴とする(2)に記載の高速延性破壊特性に優れた超高強度鋼板。 (3) The ultra-high strength steel sheet having excellent high-speed ductile fracture characteristics according to (2), wherein the average grain size of ferrite is 5 μm or less.
(4)質量%で、
C :0.03〜0.10%、
Si:0.01〜0.6%、
Mn:1.5〜2.5%、
P :0.015%以下、
S :0.003%以下、
Ni:0.1〜2.0%、
Mo:0.15〜0.60%、
Nb:0.001〜0.10%、
Ti:0.005〜0.030%、
Al:0.06%以下
を含有し、更に、
B :0.0001〜0.005%、
N :0.0001〜0.006%、
V :0.001〜0.10%、
Cu:0.01〜1.0%、
Cr:0.01〜0.8%、
Zr:0.0001〜0.005%、
Ta:0.0001〜0.005%、
Ca:0.0001〜0.01%、
REM:0.0001〜0.01%、
Mg:0.0001〜0.006%
の1種又は2種以上を含有し、残部が鉄及び不可避的不純物からなることを特徴とする(2)または(3)に記載の高速延性破壊特性に優れた超高強度鋼板。
(4) In mass%,
C: 0.03-0.10%,
Si: 0.01 to 0.6%,
Mn: 1.5 to 2.5%
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 2.0%,
Mo: 0.15-0.60%,
Nb: 0.001 to 0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less, further,
B: 0.0001 to 0.005%,
N: 0.0001 to 0.006%,
V: 0.001 to 0.10%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 0.8%
Zr: 0.0001 to 0.005%,
Ta: 0.0001 to 0.005%,
Ca: 0.0001 to 0.01%,
REM: 0.0001 to 0.01%,
Mg: 0.0001 to 0.006%
The super high strength steel sheet having excellent high-speed ductile fracture characteristics according to (2) or (3), wherein one or more of the above are contained, and the balance is iron and inevitable impurities.
(5)(1)〜(4)のいずれかに記載の鋼板の製造方法であって、(4)の成分からなる鋼を溶製、連続鋳造後、鋼片を1100〜1250℃に再加熱することを特徴とする高速延性破壊特性に優れた超高強度鋼板の製造方法。 (5) A method for producing a steel sheet according to any one of (1) to (4), wherein the steel comprising the component (4) is melted, continuously cast, and then reheated to 1100 to 1250 ° C. A method for producing an ultra-high-strength steel sheet having excellent high-speed ductile fracture characteristics.
(6)再加熱後、900℃以上の再結晶域で熱間圧延を行うことを特徴とする(5)に記載の高速延性破壊特性に優れた超高強度鋼板の製造方法。 (6) The method for producing an ultra-high strength steel sheet having excellent high-speed ductile fracture characteristics according to (5), wherein hot rolling is performed in a recrystallization region at 900 ° C. or higher after reheating.
(7)再加熱後、またはこれに引き続いた再結晶域圧延の後、880℃以下の未再結晶域で、累積圧下量が60%以上の熱間圧延を行うことを特徴とする(5)または(6)に記載の高速延性破壊特性に優れた超高強度鋼板の製造方法。 (7) It is characterized by performing hot rolling with a cumulative reduction amount of 60% or more in an unrecrystallized region of 880 ° C. or less after reheating or subsequent recrystallization region rolling (5) Or the manufacturing method of the ultra high strength steel plate excellent in the high-speed ductile fracture characteristic as described in (6).
(8)再加熱後、またはこれに引き続いた再結晶域圧延の後、880℃以下で未再結晶域圧延を開始し、600〜800℃の未再結晶域で熱間圧延を終了し、鋼板中心部の平均冷速で600〜350℃の範囲を0.5〜10℃/s以下で冷却することを特徴とする(5)〜(7)のいずれかに記載の高速延性破壊特性に優れた超高強度鋼板の製造方法。 (8) After reheating or subsequent recrystallization zone rolling, non-recrystallization zone rolling is started at 880 ° C. or less, and hot rolling is finished in the non-recrystallization zone at 600 to 800 ° C. It is excellent in the high-speed ductile fracture property according to any one of (5) to (7), characterized in that it is cooled in the range of 600 to 350 ° C at an average cold speed of the central portion at 0.5 to 10 ° C / s or less A method for manufacturing ultra-high strength steel sheets.
(9)母材が(1)〜(4)のいずれかに記載の超高強度鋼板からなることを特徴とする高速延性破壊特性に優れた超高強度鋼管。 (9) An ultra-high-strength steel pipe excellent in high-speed ductile fracture characteristics, wherein the base material is made of the ultra-high-strength steel sheet according to any one of (1) to (4).
(10)溶接金属の成分が、質量%で、
C :0.04〜0.14%、
Si:0.05〜0.4%、
Mn:1.2〜2.2%、
P :0.01%以下、
S :0.010%以下、
Ni:1.3〜3.2%、
Cr+Mo+V:1.0〜2.5%、
Ti:0.003〜0.050%、
Al:0.02%以下、
B:0.005%以下、
O:0.01〜0.03%
を含有し、残部が鉄及び不可避的不純物からなることを特徴とする(9)に記載の高速延性破壊特性に優れた超高強度鋼管。
(10) The component of the weld metal is mass%,
C: 0.04 to 0.14%,
Si: 0.05-0.4%
Mn: 1.2-2.2%,
P: 0.01% or less,
S: 0.010% or less,
Ni: 1.3-3.2%
Cr + Mo + V: 1.0 to 2.5%,
Ti: 0.003 to 0.050%,
Al: 0.02% or less,
B: 0.005% or less,
O: 0.01 to 0.03%
And the balance consists of iron and unavoidable impurities, and the super high strength steel pipe having excellent high-speed ductile fracture characteristics according to (9).
(11)(9)または(10)に記載の鋼管の製造方法であって、(1)〜(4)のいずれかに記載の超高強度鋼板をUO工程で管状に成形し、端部同士を溶接ワイヤ−及び焼成型フラックス又は溶融型フラックスを使用してサブマージドアーク溶接し、その後、拡管を行うことを特徴とする高速延性破壊特性に優れた超高強度鋼管の製造方法。 (11) A method for manufacturing a steel pipe according to (9) or (10), wherein the ultra-high strength steel sheet according to any one of (1) to (4) is formed into a tubular shape in a UO process, and ends thereof A method for producing an ultra-high strength steel pipe excellent in high-speed ductile fracture characteristics, characterized in that submerged arc welding is performed using a welding wire and a firing-type flux or a melt-type flux, followed by pipe expansion.
(12)端部同士を溶接ワイヤ−及び焼成型フラックス又は溶融型フラックスを使用してサブマージドアーク溶接を行う際に、1.5〜3.5kJ/mmの入熱にて溶接した後、拡管を行うことを特徴とする(11)に記載の高速延性破壊特性に優れた超高強度鋼管の製造方法。 (12) When submerged arc welding is performed between end portions using a welding wire and a firing-type flux or a melt-type flux, the pipe is expanded after being welded with a heat input of 1.5 to 3.5 kJ / mm. The method for producing an ultra-high strength steel pipe having excellent high-speed ductile fracture characteristics according to (11), wherein:
(13)質量%で、
C :0.01〜0.12%、
Si:0.3%以下、
Mn:1.2〜2.4%、
Ni:4.0〜8.5%、
Cr+Mo+V:3.0〜5.0%、
Ti:0.005〜0.15%、
Al:0.02%以下
を含有し、残部が鉄及び不可避的不純物からなる溶接ワイヤーを用いてサブマージドアーク溶接することを特徴とする(11)または(12)に記載の高速延性破壊特性に優れた超高強度鋼管の製造方法。
(13) In mass%,
C: 0.01 to 0.12%,
Si: 0.3% or less,
Mn: 1.2-2.4%
Ni: 4.0 to 8.5%,
Cr + Mo + V: 3.0-5.0%,
Ti: 0.005 to 0.15%,
(11) or (12), characterized in that submerged arc welding is performed using a welding wire containing Al and 0.02% or less, the balance being iron and inevitable impurities. An excellent method for manufacturing ultra-high strength steel pipes.
本発明により、高速延性破壊特性に優れた引張強さ800MPa以上(API規格X100以上)の超高強度鋼管及びその製造方法の提供が可能になり、産業上の貢献が極めて顕著である。 According to the present invention, it is possible to provide an ultra-high-strength steel pipe excellent in high-speed ductile fracture characteristics and having a tensile strength of 800 MPa or more (API standard X100 or more) and a manufacturing method thereof, and the industrial contribution is extremely remarkable.
本発明者は、鋼板の高速延性破壊特性を評価する方法について検討を行った。高速延性破壊特性は、伝播する亀裂が停止する特性であるから、亀裂の伝播のエネルギーと相関があると考えられる。そこで、種々の鋼材を用いて、シャルピー衝撃試験における荷重−変位曲線を求め、亀裂の発生のエネルギーと伝播のエネルギーを分離して評価した。その結果、800MPa以上の超高強度鋼では亀裂の発生のエネルギーが、伝播のエネルギーよりも非常に大きいことがわかった。即ち、シャルピー衝撃試験で測定した吸収エネルギーは、亀裂の発生と伝播のエネルギーを同時に評価する試験であり、亀裂の伝播のエネルギーとの相関が大きい高速延性破壊特性の評価には適さないことがわかった。これは、DWTT試験でも同様である。 This inventor examined the method of evaluating the high-speed ductile fracture characteristic of a steel plate. The high-speed ductile fracture characteristic is a characteristic in which the propagating crack stops, and is considered to have a correlation with the crack propagation energy. Therefore, using various steel materials, load-displacement curves in the Charpy impact test were obtained, and crack generation energy and propagation energy were separated and evaluated. As a result, it was found that in the ultrahigh strength steel of 800 MPa or more, the energy of crack generation is much larger than the energy of propagation. In other words, the absorbed energy measured by the Charpy impact test is a test that evaluates the energy of crack initiation and propagation at the same time, and is not suitable for the evaluation of high-speed ductile fracture characteristics that have a large correlation with the energy of crack propagation. It was. The same applies to the DWTT test.
本発明者は、次に、亀裂の伝播のエネルギーを適正に評価するための試験方法について検討を行った。板厚方向に対する板幅方向の比が大きいほど亀裂は45°方向に進展するため、DWTT試験片を用いた(フルクラックバーストテストでは亀裂は45°方向に進展している)。更に、ノッチの先端を鋭くして、亀裂の発生のエネルギーを低下させる方法として、楔状の治具に圧力を加えるプレスノッチを導入し、更に3点曲げによって延性亀裂を導入する方法を検討した。 Next, the present inventor examined a test method for properly evaluating the energy of crack propagation. Since the crack progressed in the 45 ° direction as the ratio of the plate width direction to the plate thickness direction increased, the DWTT test piece was used (in the full crack burst test, the crack progressed in the 45 ° direction). Further, as a method of reducing the energy of crack generation by sharpening the tip of the notch, a method of introducing a press notch that applies pressure to a wedge-shaped jig and further introducing a ductile crack by three-point bending was studied.
その結果、プレスノッチを試験片の中央に導入し、プレスノッチの反対側の中央部とプレスノッチ側の両端と荷重を加える3点曲げを行う際に、最大荷重に達した後、最大荷重の5%までの範囲で荷重が低下した時点で停止し、延性亀裂を導入した試験片を用いて、DWTT試験(以下、プリクラックDWTT試験という)を行えば、得られた吸収エネルギー(プリクラックDWTTエネルギーという)によって、亀裂の伝播のエネルギーを適正に評価できることがわかった。 As a result, when the press notch was introduced into the center of the test piece and the three-point bending was applied to the center part on the opposite side of the press notch and both ends on the press notch side, the maximum load was reached after reaching the maximum load. When the DWTT test (hereinafter referred to as the pre-crack DWTT test) is performed using a test piece that has been introduced with a ductile crack, it stops when the load falls within a range of up to 5%, and the obtained absorbed energy (pre-crack DWTT). It was found that the energy of propagation of cracks can be properly evaluated.
更に、プリクラックDWTT試験によって、鋼板の高速延性破壊特性が向上する要因について検討を行った。まず、0.05C−2Mn−Ni−Cu−Cr−Mo−Ti−B系の鋼板のプリクラックDWTTエネルギーとミクロ組織との関係を明確にするために検討を行った。光学顕微鏡を用いて鋼板のミクロ組織を観察し、フェライトの面積率を測定し、プリクラックDWTTエネルギーとフェライトの面積率との関係を調査した。その結果、図1に示すように、プリクラックDWTTエネルギーは、ミクロ組織のフェライトの面積率と相関があることがわかった。 Furthermore, the factor which improves the high-speed ductile fracture characteristic of a steel plate was examined by the precrack DWTT test. First, in order to clarify the relationship between the precrack DWTT energy and the microstructure of a 0.05C-2Mn-Ni-Cu-Cr-Mo-Ti-B steel sheet. The microstructure of the steel sheet was observed using an optical microscope, the area ratio of ferrite was measured, and the relationship between the precrack DWTT energy and the area ratio of ferrite was investigated. As a result, as shown in FIG. 1, it was found that the precrack DWTT energy has a correlation with the area ratio of ferrite in the microstructure.
更に、本発明者は、プリクラックDWTTエネルギーと集合組織の関係を詳細に調査した。その結果、図2に示すように、圧延方向から45°方向の板厚断面での(100)面の集積が3以上になるとプリクラックDWTTエネルギーが著しく低下することがわかった。なお、(100)面の集積は、X線回折による強度の測定値が、ランダムな方位を有する標準試料の何倍であるかによって評価する。即ち、(100)面の集積が3とは、ランダムな方位を有する標準試料に比べて、(100)面の強度が3倍であることを意味する。 Furthermore, the inventor has investigated in detail the relationship between the precrack DWTT energy and the texture. As a result, as shown in FIG. 2, it was found that the precrack DWTT energy was remarkably reduced when the accumulation of (100) planes in the plate thickness section in the 45 ° direction from the rolling direction was 3 or more. The (100) plane accumulation is evaluated according to how many times the measured value of the intensity by X-ray diffraction is a standard sample having a random orientation. That is, the accumulation of the (100) plane of 3 means that the intensity of the (100) plane is three times that of the standard sample having a random orientation.
圧延方向を軸として圧延面から45°回転した面の断面(この面の法線方向をD方向という)の(100)面の集積とプリクラックDWTTエネルギーに相関が認められる理由は、延性破壊の亀裂が圧延方向から45°の方向に進展し、この方向に鋼の劈開面に相当する(100)面が集積すると、亀裂が一気に伝播するためであると考えられる。また、D方向の板厚断面で、(100)面が集積した部位のミクロ組織を観察した結果、主にベイナイト・マルテンサイトであることがわかった。即ち、熱間圧延後のオーステナイトはD方向の板厚断面に(100)が集積している。一方、熱間圧延後のフェライトは圧延方向の板厚断面に(100)が集積している。したがって、熱間圧延後のフェライト分率を増加させれば、D方向の板厚断面への(100)の集積度が低下し、延性亀裂の伝播を抑制できる。 The reason why the accumulation of the (100) plane of the cross section of the plane rotated 45 ° from the rolling plane with the rolling direction as the axis (the normal direction of this plane is called the D direction) and the precrack DWTT energy is recognized is that of ductile fracture It is considered that the crack propagates all at once when the crack propagates in the direction of 45 ° from the rolling direction and the (100) plane corresponding to the cleavage plane of steel accumulates in this direction. Moreover, as a result of observing the microstructure of the part where (100) planes were accumulated in the plate thickness cross section in the D direction, it was found that it was mainly bainite martensite. That is, austenite after hot rolling has (100) accumulated in the plate thickness section in the D direction. On the other hand, in the ferrite after hot rolling, (100) is accumulated on the plate thickness section in the rolling direction. Therefore, if the ferrite fraction after hot rolling is increased, the degree of accumulation of (100) in the thickness direction cross section in the D direction is reduced, and the propagation of ductile cracks can be suppressed.
また、このような鋼板を製造する方法について詳細に検討を行った。ミクロ組織がフェライトとベイナイト・マルテンサイトの2相からなる800MPa以上の引張強度を有する高強度鋼を製造する際に、熱間圧延の仕上げ温度をAr3点以下とし、水冷するとフェライトが生成する。本発明者は、0.05C−2Mn−Ni−Cu−Cr−Mo−Ti−B系の鋼板を種々の条件で製造し、熱間圧延の仕上温度とフェライトの面積率の関係を調査した。その結果、図3に示したように、熱間圧延の仕上温度とフェライトの面積率には相関があることがわかった。 Moreover, it examined in detail about the method of manufacturing such a steel plate. When producing a high strength steel having a tensile strength of 800 MPa or more, the microstructure of which is composed of two phases of ferrite and bainite / martensite, the hot rolling finish temperature is set to Ar 3 point or less, and ferrite is produced when cooled with water. This inventor manufactured the 0.05C-2Mn-Ni-Cu-Cr-Mo-Ti-B type steel plate on various conditions, and investigated the relationship between the finishing temperature of hot rolling, and the area ratio of a ferrite. As a result, as shown in FIG. 3, it was found that there is a correlation between the hot rolling finishing temperature and the area ratio of ferrite.
以下、本発明について詳細に説明する。 Hereinafter, the present invention will be described in detail.
フェライトの面積率は、1%未満ではD方向の板厚断面に(100)面が集積したベイナイト・マルテンサイトが多くなり、耐高速延性破壊特性が低下する。一方、フェライトの面積率が60%超では、強度が低下する。強度と耐高速延性破壊特性のバランスを考慮すると、フェライトの面積率は、好ましくは5%超〜20%以下である。残部はベイナイト・マルテンサイトである。しかし、フェライトの面積率が1〜5%の範囲であっても必要な強度と耐高速延性破壊特性は確保できる。 When the area ratio of the ferrite is less than 1%, bainite martensite having (100) planes accumulated in the thickness direction cross section in the D direction increases, and the high-speed ductile fracture resistance deteriorates. On the other hand, when the area ratio of ferrite exceeds 60%, the strength decreases. Considering the balance between strength and high-speed ductile fracture resistance, the area ratio of ferrite is preferably more than 5% to 20%. The balance is bainite martensite. However, even if the ferrite area ratio is in the range of 1 to 5%, the necessary strength and high-speed ductile fracture resistance can be ensured.
ミクロ組織において、フェライトとベイナイト・マルテンサイトは、光学顕微鏡あるいは走査電子顕微鏡を用いた組織観察によって判別することが可能であり、光学顕微鏡あるいは走査電子顕微鏡で撮影した組織写真を画像解析することにより、フェライトの面積率を測定することができる。 In the microstructure, ferrite and bainite martensite can be distinguished by structural observation using an optical microscope or scanning electron microscope, and by analyzing the structure photograph taken with an optical microscope or scanning electron microscope, The area ratio of ferrite can be measured.
D方向の板厚断面での(100)面の集積は、3を超えると高速延性破壊特性が低下するため、3以下とした。なお、D方向の板厚断面での(100)面の集積は、ランダムな方位を有する標準試料のX線回折による強度の測定値を基準とするものであるから、下限は1である。 The integration of (100) planes in the cross section of the plate thickness in the D direction is set to 3 or less because if it exceeds 3, the high-speed ductile fracture characteristics deteriorate. Note that the accumulation of the (100) plane in the thickness direction cross section in the D direction is based on the measured value of the intensity by X-ray diffraction of a standard sample having a random orientation, so the lower limit is 1.
また、フェライトの平均粒径が5μmを超えて粗大化すると、劈開の破面単位が大きくなって伝播エネルギーが低下することがあるため、フェライトの平均粒径は、5μm以下であることが好ましい。フェライトの平均粒径は、光学顕微鏡又は走査型電子顕微鏡によって撮影した組織写真を用いて、切断法によって測定することができる。 Further, if the average grain size of ferrite exceeds 5 μm, the cleaved fracture surface unit becomes large and the propagation energy may decrease, so the average grain size of ferrite is preferably 5 μm or less. The average particle diameter of ferrite can be measured by a cutting method using a structure photograph taken with an optical microscope or a scanning electron microscope.
次に母材の化学成分の限定理由について説明する。 Next, the reason for limiting the chemical component of the base material will be described.
Cは鋼の強度向上に極めて有効な元素であり、0.03%以上のCを含有することが好ましい。しかし、C含有量が0.10%よりも多すぎると母材及び溶接熱影響部(HAZという)の低温靱性がやや劣化し、現地溶接性を損なうことがあるため、C含有量の上限を0.10%以下とすることが好ましく、0.07%以下とすることが更に好ましい。 C is an extremely effective element for improving the strength of steel, and preferably contains 0.03% or more of C. However, if the C content is more than 0.10%, the low temperature toughness of the base metal and the weld heat affected zone (referred to as HAZ) may be slightly deteriorated, which may impair on-site weldability. The content is preferably 0.10% or less, and more preferably 0.07% or less.
Siは、脱酸に有効な元素であり、0.01%以上を含有することが好ましいが、0.6%よりも多く添加するとHAZの低温靱性がやや劣化し、現地溶接性を損なうことがあるため、Si含有量の上限を0.6%以下とすることが好ましい。 Si is an element effective for deoxidation, and it is preferable to contain 0.01% or more. However, if it is added more than 0.6%, the low temperature toughness of HAZ is slightly deteriorated and the on-site weldability may be impaired. For this reason, the upper limit of the Si content is preferably 0.6% or less.
Mnは、鋼の強度と低温靱性とのバランスを良好にするために有効な元素であり、Mn含有量の下限を1.5%以上とすることが好ましい。しかし、Mnを2.5%よりも過剰に含有すると鋼の焼き入れ性が増してHAZの低温靱性を劣化させ、また、現地溶接性を損なうことがある。したがって、Mn含有量の上限を2.5%以下とすることが好ましい。 Mn is an element effective for improving the balance between strength and low temperature toughness of steel, and the lower limit of the Mn content is preferably 1.5% or more. However, if Mn is contained in excess of 2.5%, the hardenability of the steel is increased and the low temperature toughness of the HAZ is deteriorated, and the on-site weldability may be impaired. Therefore, it is preferable that the upper limit of the Mn content is 2.5% or less.
P、Sは不純物元素であり、母材及びHAZの低温靱性をより一層向上させるために、Pの含有量及びSの含有量の上限をそれぞれ0.015%以下及び0.003%以下とすることが好ましい。Pの含有量及びSの含有量の下限は低いほど好ましいため規定しないが、通常、それぞれ0.001%以上及び0.0001%以上を含有する。 P and S are impurity elements, and in order to further improve the low temperature toughness of the base material and the HAZ, the upper limits of the P content and the S content are 0.015% or less and 0.003% or less, respectively. It is preferable. Although the lower limit of the P content and the S content is preferably as low as possible, it is not specified, but usually contains 0.001% or more and 0.0001% or more, respectively.
Niは、低温靱性及び強度を向上させる元素であり、その効果を得るために、Ni含有量の下限を0.1%以上とすることが好ましい。一方、Niの含有量が2.0%を超えると、溶接性を損なうことがあるため、Ni含有量の上限を2.0%とすることが好ましい。 Ni is an element that improves low-temperature toughness and strength, and in order to obtain the effect, the lower limit of the Ni content is preferably set to 0.1% or more. On the other hand, if the Ni content exceeds 2.0%, weldability may be impaired, so the upper limit of the Ni content is preferably 2.0%.
Moは、鋼の焼き入れ性を向上させ、炭窒化物を形成して強度を向上させる元素であり、その効果を得るには、Mo含有量を0.15%以上とすることが好ましい。一方、Moを0.60%超含有すると、強度が高くなり過ぎてHAZの低温靱性を損なうことがあるため、Mo含有量の上限を0.60%とすることが好ましい。 Mo is an element that improves the hardenability of steel and forms carbonitride to improve strength. To obtain the effect, Mo content is preferably 0.15% or more. On the other hand, if the Mo content exceeds 0.60%, the strength becomes too high and the low temperature toughness of the HAZ may be impaired, so the upper limit of the Mo content is preferably 0.60%.
Nbは炭化物、窒化物を形成し、鋼の強度を向上させる元素であり、この効果を得るには、Nb含有量を0.001%以上とすることが好ましい。一方、Nb含有量が0.10%よりも多すぎると、母材及びHAZの低温靱性を損なうことがあるため、Nb含有量の上限を0.10%とすることが好ましい。 Nb is an element that forms carbides and nitrides and improves the strength of the steel. To obtain this effect, the Nb content is preferably 0.001% or more. On the other hand, if the Nb content is more than 0.10%, the low temperature toughness of the base material and the HAZ may be impaired, so the upper limit of the Nb content is preferably 0.10%.
Tiは、脱酸に有効であり、窒化物を形成して結晶粒径の微細化に寄与する元素であり、その効果を得るには、0.005%以上を添加することが好ましい。一方、Ti含有量が0.030%よりも多すぎると、粗大な炭化物を生じて、低温靱性を劣化させることがあるため、Ti含有量の上限を0.030%以下とすることが好ましい。 Ti is an element that is effective for deoxidation and contributes to the refinement of the crystal grain size by forming a nitride. To obtain the effect, it is preferable to add 0.005% or more. On the other hand, if the Ti content is more than 0.030%, coarse carbides may be produced and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably 0.030% or less.
Alは脱酸材として有効な元素であるが、Al含有量が0.06%を超えるとAl系非金属介在物が増加して鋼の清浄度を阻害することがあるため、Al含有量の上限を0.06%以下とした。また、脱酸はTi及び/又はSiでも可能であるため、Alを必ずしも含有する必要はなく、下限は0%でも良い。 Al is an effective element as a deoxidizing material, but if the Al content exceeds 0.06%, Al-based non-metallic inclusions may increase and inhibit the cleanliness of the steel. The upper limit was made 0.06% or less. Further, since deoxidation can be performed with Ti and / or Si, it is not always necessary to contain Al, and the lower limit may be 0%.
なお、本発明においては、強度および靱性を改善する元素として、B、N、V、Cu、Cr、Zr、Ta、Ca,REM、Mgの1種または2種以上の元素を添加することができる。 In the present invention, one or more elements of B, N, V, Cu, Cr, Zr, Ta, Ca, REM, and Mg can be added as elements for improving strength and toughness. .
Bは、焼入れ性を高め、溶接熱影響部の靱性を向上させる元素である。この効果は、0.0001%以上の添加で顕著になるが、0.005%よりも過剰の添加は、靱性の低下を招くことがある。したがって、Bの添加量を0.0001〜0.005%の範囲とすることが好ましい。 B is an element that enhances hardenability and improves the toughness of the weld heat affected zone. This effect becomes prominent with addition of 0.0001% or more, but addition exceeding 0.005% may lead to a decrease in toughness. Therefore, it is preferable that the addition amount of B is in the range of 0.0001 to 0.005%.
Nは、Ti、Al等と窒化物を形成し、溶接熱影響部のオーステナイト粒の粗大化を防止する。この効果は、0.0001%以上の添加で顕著になるが、0.006%よりも過剰の添加は、靱性の低下を招くことがある。したがって、Nの添加量を0.0001〜0.006%の範囲とすることが好ましい。 N forms nitrides with Ti, Al, etc., and prevents the austenite grains in the weld heat affected zone from becoming coarse. This effect becomes prominent with addition of 0.0001% or more, but addition exceeding 0.006% may lead to a decrease in toughness. Therefore, it is preferable that the addition amount of N is in the range of 0.0001 to 0.006%.
Vは、Nbと同様に炭化物、窒化物を形成し、鋼の強度を向上させる元素であるが、顕著な効果を得るには0.01%以上の添加が好ましい。一方、Vを0.10%超添加すると、靱性の低下を招くことがあるため、上限を0.10%以下とすることが好ましい。 V is an element that forms carbides and nitrides in the same manner as Nb and improves the strength of the steel. However, in order to obtain a remarkable effect, V is preferably added in an amount of 0.01% or more. On the other hand, if V is added in excess of 0.10%, the toughness may be lowered, so the upper limit is preferably made 0.10% or less.
Cuは、強度を上昇させる元素であり、0.01%以上添加することが好ましい。一方、1.0%超を添加すると鋼片加熱時や溶接時に割れを生じやすくするため、上限を1.0%以下とすることが好ましい。 Cu is an element that increases the strength, and is preferably added in an amount of 0.01% or more. On the other hand, if more than 1.0% is added, cracking is likely to occur during heating of the steel slab or during welding, so the upper limit is preferably made 1.0% or less.
Crは、析出強化によって鋼の強度を向上させる元素であり、0.01%以上の添加が有効である。一方、0.8%よりも多量に添加すると、鋼の焼入れ性を上昇させて、靱性を低下させることがあるため、上限を0.8%以下とすることが好ましい。 Cr is an element that improves the strength of steel by precipitation strengthening, and the addition of 0.01% or more is effective. On the other hand, if added in a larger amount than 0.8%, the hardenability of the steel is increased and the toughness may be lowered, so the upper limit is preferably made 0.8% or less.
Zr及びTaは、Nbと同様に炭化物、窒化物を形成し、鋼の強度を向上させる元素であり、それぞれ、0.0001%以上の添加が好ましい。一方、Zr及びTaを、それぞれ、0.0050%超添加すると、靱性の低下を招くことがある。そのため、Zr及びTaの添加量の上限をそれぞれ、0.005%以下とすることが好ましい。 Zr and Ta are elements that form carbides and nitrides in the same manner as Nb and improve the strength of the steel, and are each preferably added in an amount of 0.0001% or more. On the other hand, if Zr and Ta are added in excess of 0.0050%, toughness may be reduced. Therefore, it is preferable that the upper limit of the addition amount of Zr and Ta is 0.005% or less, respectively.
Ca及びREMは硫化物を生成することにより、伸長したMnSの生成を抑制し、鋼材の板厚方向の特性、特に耐ラメラティアー性を改善する。この効果を得るには、Ca及びREMを、それぞれ、0.0001%以上添加することが好ましい。一方、Ca及びREMを、それぞれ、0.01%超添加すると、Ca及びREMの酸化物が増加する。そのため、Ca及びREMの添加量の上限を、それぞれ、0.01%以下とすることが好ましい。 Ca and REM suppress the production | generation of the extended | stretched MnS by producing | generating a sulfide, and improve the characteristic in the plate | board thickness direction of steel materials, especially the lamellar resistance. In order to obtain this effect, it is preferable to add 0.0001% or more of Ca and REM, respectively. On the other hand, if Ca and REM are added in excess of 0.01%, Ca and REM oxides increase. Therefore, it is preferable that the upper limit of the addition amount of Ca and REM is 0.01% or less, respectively.
Mgは、MgO、MgS等の超微細なMg含有酸化物又は硫化物を生成し、オーステナイト粒の粗大化を抑制し、HAZ靱性を向上させる元素である。この効果を得るには、Mgを0.0001%以上添加することが好ましい。一方、Mgを0.006%超添加するとMg含有酸化物、硫化物が粗大化するため、その上限を0.006%以下とすることが好ましい。 Mg is an element that generates ultrafine Mg-containing oxides or sulfides such as MgO and MgS, suppresses coarsening of austenite grains, and improves HAZ toughness. In order to acquire this effect, it is preferable to add 0.0001% or more of Mg. On the other hand, if Mg is added in excess of 0.006%, the Mg-containing oxide and sulfide are coarsened, so the upper limit is preferably made 0.006% or less.
上記の鋼板を鋼管とする場合の、溶接金属の成分の限定理由について述べる。 The reason for limiting the components of the weld metal when the steel plate is a steel pipe will be described.
Cは、鋼の強度向上に極めて有効であり、マルテンサイト組織において目標とする強度を得るためには、C含有量を0.04%以上とすることが好ましい。一方、C含有量が0.14%を超えると溶接低温割れが発生しやすくなり、現地溶接部とシーム溶接が交わる、いわゆるTクロス部のHAZ最高硬さの上昇を招くので、C含有量の上限を0.14%以下とすることが好ましい。更に好ましいC含有量の上限値は0.10%以下である。 C is extremely effective for improving the strength of the steel, and in order to obtain the target strength in the martensite structure, the C content is preferably 0.04% or more. On the other hand, if the C content exceeds 0.14%, welding low temperature cracking is likely to occur, and the on-site welded part and seam welding intersect, leading to an increase in the HAZ maximum hardness of the so-called T-cross part. The upper limit is preferably 0.14% or less. A more preferable upper limit of the C content is 0.10% or less.
Siは、ブローホールの発生を防止するために、0.05%以上含有させることが好ましい。一方、Si含有量が0.4%よりも多いと、低温靱性を劣化させることがあり、特に、内外面溶接や多層溶接を行う場合、再熱部の低温靱性を劣化させることがあるため、上限を0.4%以下とすることが好ましい。 In order to prevent the occurrence of blow holes, Si is preferably contained in an amount of 0.05% or more. On the other hand, when the Si content is more than 0.4%, the low temperature toughness may be deteriorated, and particularly when performing inner and outer surface welding or multilayer welding, the low temperature toughness of the reheated portion may be deteriorated. The upper limit is preferably 0.4% or less.
Mnは、強度、低温靱性のバランスを良好にし、粒内ベイナイトの生成核となる介在物を形成する元素である。この効果を得るには、Mn含有量を1.2%以上にすることが好ましい。一方、Mn含有量が2.2%よりも多すぎると偏析が助長され、低温靱性が劣化することがあり、溶接材料の製造が困難になるので、Mn含有量の上限を2.2%以下とすることが好ましい。 Mn is an element that improves the balance between strength and low-temperature toughness and forms inclusions that form nuclei for intragranular bainite. In order to obtain this effect, the Mn content is preferably 1.2% or more. On the other hand, if the Mn content is more than 2.2%, segregation is promoted and the low temperature toughness may be deteriorated, making it difficult to produce a welding material. Therefore, the upper limit of the Mn content is 2.2% or less. It is preferable that
P、Sは不可避的不純物であり、低温靱性の劣化を抑制し、低温割れ感受性を低減するためには、少ないほど好ましく、P、Sの含有量を、それぞれ、0.001%以下、0.001%以下とすることが好ましい。 P and S are inevitable impurities, and in order to suppress the low temperature toughness deterioration and reduce the low temperature cracking susceptibility, the smaller the content, the smaller the content of P and S, 0.001% or less, respectively. It is preferable to be 001% or less.
Niは、焼き入れ性を高めて強度を向上させ、低温靱性を向上させる元素であり、この効果を得るためには、1.3%以上のNiを含有させることが好ましい。一方、Ni含有量が3.2%よりも多すぎると高温割れを生じることがあるため、Ni含有量の上限を3.2%以下とすることが好ましい。 Ni is an element that enhances hardenability and improves strength and improves low-temperature toughness. In order to obtain this effect, it is preferable to contain 1.3% or more of Ni. On the other hand, if the Ni content is more than 3.2%, hot cracking may occur, so the upper limit of the Ni content is preferably 3.2% or less.
Cr、Mo、Vは、何れも焼き入れ性を高め、強度を向上させる元素であり、効果を得るには、Cr+Mo+Vを1.0%以上とすることが好ましい。一方、Cr+Mo+Vを2.5%よりも多量に添加すると低温割れを生じることがあるため、Cr+Mo+V含有量の上限を2.5%以下とすることが好ましい。 Cr, Mo, and V are all elements that increase the hardenability and improve the strength. To obtain the effect, Cr + Mo + V is preferably set to 1.0% or more. On the other hand, if Cr + Mo + V is added in a larger amount than 2.5%, low temperature cracking may occur, so the upper limit of the Cr + Mo + V content is preferably 2.5% or less.
Tiは、粒内ベイナイトの生成核となるTiの窒化物及び酸化物等を形成する元素であり、0.003%以上を含有させることが好ましい。一方、Ti含有量が0.05%よりも多すぎると、Tiの炭化物が多く生成し、低温靱性を劣化させることがあるため、Ti含有量の上限を0.05%とすることが好ましい。 Ti is an element that forms a nitride, oxide, or the like of Ti that serves as a nucleus for formation of intragranular bainite, and preferably contains 0.003% or more. On the other hand, when the Ti content is more than 0.05%, a large amount of Ti carbide is generated and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably 0.05%.
Alは、粒内ベイナイトの生成核となるTiの酸化物の生成を阻害することがあるため、Al含有量は少ない方が好ましい。Al含有量の好ましい上限は0.02%以下であり、更に好ましくは0.015%以下が良い。 Since Al sometimes inhibits the formation of Ti oxides that form nuclei of intragranular bainite, it is preferable that the Al content is low. The upper limit with preferable Al content is 0.02% or less, More preferably, 0.015% or less is good.
Bは、焼き入れ性を高め、溶接金属の低温靱性を向上させる元素であり、0.0003%以上を含有することが好ましいが、B含有量が0.005%よりも多すぎると低温靱性を劣化させることがあるため、B含有量を0.005%以下とすることが好ましい。 B is an element that enhances the hardenability and improves the low temperature toughness of the weld metal, and preferably contains 0.0003% or more, but if the B content is more than 0.005%, the low temperature toughness is reduced. Since it may deteriorate, the B content is preferably 0.005% or less.
Oは、焼入れ性を下げ、溶接金属の低温靭性を劣化させる元素であり、O量が0.03%を超えると低温靭性を著しく劣化させる。一方、O量が低いと低温割れが発生しやすくなると同時に現地溶接性の硬さが高くなるので0.010%以上とした。 O is an element that lowers the hardenability and degrades the low temperature toughness of the weld metal. When the amount of O exceeds 0.03%, the low temperature toughness is remarkably deteriorated. On the other hand, if the amount of O is low, cold cracking is likely to occur, and at the same time, the hardness of on-site weldability increases, so the content was made 0.010% or more.
溶接金属には、その他に溶接時の精錬・凝固を良好に行わせるために添加させたZr、Nb、Mg等の元素を含有する場合がある。 In addition, the weld metal may contain elements such as Zr, Nb, and Mg that are added to improve the refining and solidification during welding.
溶接金属の組織は、主にベイナイト・マルテンサイト、粒内ベイナイトからなり、残部はフェライト及び/又は残留オーステナイトである。引張強度を800MPa以上にするために、ベイナイト・マルテンサイトの面積率を50%以上にすることが好ましい。 The structure of the weld metal is mainly composed of bainite / martensite and intragranular bainite, and the balance is ferrite and / or retained austenite. In order to set the tensile strength to 800 MPa or more, it is preferable to set the area ratio of bainite / martensite to 50% or more.
更に溶接金属の低温靱性を良好にするには粒内ベイナイトの面積率が多ければ多い方が好ましく、10%以上にした方がよい。ベイナイト・マルテンサイトと粒内ベイナイトは、光学顕微鏡又は走査型電子顕微鏡による組織観察によって判別することができ、ベイナイト・マルテンサイト、粒内ベイナイトの面積率は、光学顕微鏡又は走査型電子顕微鏡によって撮影した組織写真を用いて画像解析によって測定することができる。 Furthermore, in order to improve the low temperature toughness of the weld metal, it is preferable that the area ratio of intragranular bainite is as large as possible. Bainite martensite and intragranular bainite can be distinguished by observation of the structure with an optical microscope or scanning electron microscope, and the area ratio of bainite martensite and intragranular bainite was photographed with an optical microscope or scanning electron microscope. It can be measured by image analysis using tissue photographs.
次に高速延性破壊特性を良好にする鋼板の製造方法について説明する。上記に示した成分を含有する鋼を製鋼工程で溶製後、連続鋳造し、その後、加熱し、熱間圧延を施す。鋼片の加熱温度は1100〜1250℃に規定する。1100℃未満では粗大なγ粒が存在し、鋼板までその粗大粒がのこる。このため加熱温度を1100℃以上とした。一方1250℃を越えると粒成長が起こるためこれまた粗大粒が生成しやすくなり、低温靭性を劣化させるので1250℃以下にした。
次に再結晶域圧延について述べる。再結晶域圧延の圧延温度が900℃以上とした。再結晶域圧延の圧延温度が900℃未満になると、オーステナイトの十分な再結晶化が図れず、結晶粒が細粒化しないため再結晶域圧延の圧延温度を900℃以上とした。
Next, the manufacturing method of the steel plate which makes a high-speed ductile fracture characteristic favorable is demonstrated. The steel containing the components shown above is melted in the steel making process, continuously cast, then heated, and hot rolled. The heating temperature of the steel slab is specified at 1100 to 1250 ° C. When the temperature is lower than 1100 ° C., coarse γ grains exist, and the coarse grains remain up to the steel sheet. For this reason, heating temperature was 1100 degreeC or more. On the other hand, when the temperature exceeds 1250 ° C., grain growth occurs, so that coarse grains are likely to be formed, and the low temperature toughness is deteriorated.
Next, recrystallization zone rolling will be described. The rolling temperature of the recrystallization zone rolling was set to 900 ° C. or higher. When the rolling temperature of the recrystallization zone rolling is less than 900 ° C., the austenite cannot be sufficiently recrystallized, and the crystal grains do not become fine. Therefore, the rolling temperature of the recrystallization zone rolling is set to 900 ° C. or higher.
次に、未再結晶域圧延の条件について説明する。
本発明の未再結晶域での圧延条件は、圧延温度を880℃以下とし、かつ、累積圧下率を60%以上とする。880℃を超えると板厚中心部では一部再結晶域圧延になってしまうために粒の細粒化がはかれないので圧下温度を880℃以下とした。また、累積圧下量を60%未満では結晶粒径が微細化しないため60%以上とした。
Next, the conditions for non-recrystallization zone rolling will be described.
The rolling conditions in the non-recrystallized region of the present invention are a rolling temperature of 880 ° C. or lower and a cumulative rolling reduction of 60% or higher. If the temperature exceeds 880 ° C., the recrystallization zone rolling is partially performed at the center of the plate thickness, so that the grain size cannot be reduced. In addition, when the cumulative reduction amount is less than 60%, the crystal grain size is not refined, so that it is set to 60% or more.
一方、フェライトの面積率を上げるには、未再結晶域での熱間圧延の終了温度を低下させると増加する。フェライトの面積率を1〜60%とするには、未再結晶域での熱間圧延の終了温度を800℃以下にすることが必要であり、800℃以下とした。好ましい範囲は、600〜780℃である。 On the other hand, in order to increase the area ratio of ferrite, it increases when the end temperature of hot rolling in the non-recrystallized region is lowered. In order to set the area ratio of ferrite to 1 to 60%, it is necessary to set the end temperature of hot rolling in the non-recrystallized region to 800 ° C. or less, and to 800 ° C. or less. A preferred range is 600-780 ° C.
さらに、熱間圧延の終了後、600〜350℃の範囲を0.1〜10℃/sで冷却する。この冷却は制御しやすい水冷が望ましい。冷却速度が0.1℃/s未満では、未再結晶域圧延の終了時には微細であったオーステナイト粒が成長し、平均旧オーステナイト粒径が5μm超となり、低温靭性が低下する。この冷却速度はオーステナイト粒成長を避けるために1℃/s以上であることが好ましい。また、10℃/s超では鋼板表面近傍のフェライトの面積率が1%未満となる場合があり、10℃/s以下である必要がある。 Furthermore, after completion | finish of hot rolling, the range of 600-350 degreeC is cooled at 0.1-10 degreeC / s. This cooling is preferably water cooling which is easy to control. When the cooling rate is less than 0.1 ° C./s, fine austenite grains grow at the end of the non-recrystallization zone rolling, the average prior austenite grain size exceeds 5 μm, and the low temperature toughness decreases. This cooling rate is preferably 1 ° C./s or more in order to avoid austenite grain growth. If it exceeds 10 ° C./s, the area ratio of ferrite near the surface of the steel sheet may be less than 1%, and needs to be 10 ° C./s or less.
更に、鋼板を筒状にプレス成形し、端部同士をサブマージアーク溶接して鋼管とする。 Furthermore, the steel plate is press-formed into a cylindrical shape, and the ends are submerged arc welded to form a steel pipe.
サブマージアーク溶接は母材の希釈が大きい溶接であり、所望の特性すなわち溶接金属組成を得るためには、母材の希釈を考慮した溶接材料の選択が必要である。以下、溶接ワイヤーの化学組成の限定理由を述べるが、基本的には超高強度ラインパイプを実現できる製造方法である。 Submerged arc welding is a welding with a large dilution of the base metal, and in order to obtain a desired characteristic, that is, a weld metal composition, it is necessary to select a welding material in consideration of the dilution of the base metal. Hereinafter, although the reason for limiting the chemical composition of the welding wire will be described, it is basically a manufacturing method capable of realizing an ultra-high strength line pipe.
Cは、溶接金属で必要とされる範囲のC含有量を得るために、母材成分による希釈及び雰囲気からCの混入を考慮して0.01〜0.12%とした。 In order to obtain the C content in a range required for the weld metal, C is set to 0.01 to 0.12% in consideration of dilution by the base material component and mixing of C from the atmosphere.
Si、Mn、Ni、Cr+Mo+Vは、溶接金属で必要とされる範囲のSi、Mn、Ni、Cr+Mo+Vの含有量を得るために、母材成分による希釈を考慮して、それぞれ、0.3%以下、1.2〜2.4%、4.0〜8.5%、3.0〜5.0%とした。 Si, Mn, Ni, Cr + Mo + V is 0.3% or less in consideration of dilution by the base material component in order to obtain the content of Si, Mn, Ni, Cr + Mo + V in the range required for the weld metal. 1.2-2.4%, 4.0-8.5%, 3.0-5.0%.
Tiは、粒内ベイナイトの生成核となるTiの窒化物及び酸化物等を形成する元素であり、0.005%以上を含有させることが好ましい。一方、Ti含有量が0.15%よりも多すぎると、Tiの炭化物が多く生成し、低温靱性を劣化させることがあるため、Ti含有量の上限を0.15%とすることが好ましい。 Ti is an element that forms a nitride, oxide, or the like of Ti that forms nuclei for intragranular bainite, and it is preferable to contain 0.005% or more. On the other hand, if the Ti content is more than 0.15%, a large amount of Ti carbide is generated and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably set to 0.15%.
Alは、粒内ベイナイトの生成核となるTiの酸化物の生成を阻害することがあるため、Al含有量は少ない方が好ましい。Al含有量の好ましい上限は0.02%以下である。 Since Al sometimes inhibits the formation of Ti oxides that form nuclei of intragranular bainite, it is preferable that the Al content is low. The upper limit with preferable Al content is 0.02% or less.
その他P,Sの不純物は極力少ない方が望ましく、Bは強度確保に添加することも可能である。また、Zr,Nb,Mg等が脱酸を目的として使用される。 In addition, it is desirable that impurities of P and S are as small as possible, and B can be added to ensure strength. Zr, Nb, Mg, etc. are used for the purpose of deoxidation.
なお、溶接は単極だけでなく、複数電極での溶接も可能である。複数電極で溶接の場合は各種ワイヤーの組み合わせが可能であり、個々のワイヤーが上記成分範囲にある必要はなく、それぞれのワイヤー成分と消費量からの平均組成が上記成分範囲にあれば良い。 In addition, welding can be performed with a plurality of electrodes as well as a single electrode. In the case of welding with a plurality of electrodes, it is possible to combine various wires, and it is not necessary for each wire to be in the above-mentioned component range, and the average composition from each wire component and consumption may be in the above-described component range.
サブマージドアーク溶接に使用されるフラックスは大別すると焼成型フラックスと溶融型フラックスがある。焼成型フラックスは合金材添加が可能で拡散性水素量が低い利点があるが、粉化しやすく繰り返し使用が難しい欠点がある。一方、溶融型フラックスはガラス粉状で、粒強度が高く、吸湿しにくい利点があり、拡散性水素がやや高い欠点がある。本発明の超高強度鋼管を製造する場合には、溶接低温割れが起こりやすく、この点からは焼成型が望ましいが、一方、回収して繰り返し使用が可能な溶融型は大量生産に向きコストが低い利点がある。焼成型ではコストが高いことが、溶融型では厳密な品質管理の必要性が問題であるが、工業的に対処可能な範囲であり、どちらでも本質的には使用可能である。 Flux used for submerged arc welding can be broadly classified into fired flux and molten flux. Firing-type fluxes have the advantage that alloy materials can be added and the amount of diffusible hydrogen is low, but they have the disadvantage of being easily pulverized and difficult to use repeatedly. On the other hand, the melt-type flux is in the form of glass powder, has the advantages of high grain strength and is difficult to absorb moisture, and has the disadvantage that diffusible hydrogen is slightly high. When manufacturing the ultra-high strength steel pipe of the present invention, cold cracking is likely to occur. From this point, a fired mold is desirable, but a molten mold that can be recovered and used repeatedly is suitable for mass production and has a low cost. There are low advantages. The cost is high in the baking mold, and the necessity of strict quality control is a problem in the melting mold, but it is within a range that can be handled industrially, and either can be used essentially.
次に溶接条件について以下に説明する。 Next, welding conditions will be described below.
最初に行う仮付け溶接は、MAGアーク溶接、MIGアーク溶接、TIGアーク溶接の何れでもよい。通常はMAGアーク溶接である。次に内外面の溶接を、サブマージドアーク溶接とすることが好ましいが、TIGアーク溶接、MIGアーク溶接、MAGアーク溶接でも良い。内外面の溶接はそれぞれ1パスづつでも良いが、複数パス行っても良い。 The initial tack welding performed may be any of MAG arc welding, MIG arc welding, and TIG arc welding. Usually, MAG arc welding. Next, the inner and outer surfaces are preferably submerged arc welding, but may be TIG arc welding, MIG arc welding, or MAG arc welding. The inner and outer surfaces may be welded one by one, but a plurality of passes may be performed.
内外面をサブマージドアーク溶接する場合、溶接速度を1m/分未満とするとラインパイプのシーム溶接としては非効率であり、3m/分を超えるとビード形状が不安定になることがある。したがって、サブマージドアーク溶接の溶接速度は、1〜3m/分の範囲内であることが好ましい。 When submerged arc welding is performed on the inner and outer surfaces, if the welding speed is less than 1 m / min, seam welding of the line pipe is inefficient, and if it exceeds 3 m / min, the bead shape may become unstable. Therefore, the welding speed of submerged arc welding is preferably in the range of 1 to 3 m / min.
なお、仮付け溶接と内外面の溶接の溶接部が重複する場合には、溶接入熱は出来る限り低い方が好ましい。また、溶接入熱は板厚によって異なるが、入熱が小さすぎると溶け込みが不十分になり、溶接回数が多くなり、作業効率が悪くなり、溶接入熱が大きすぎると熱影響部の軟化が大きく、溶接部の靭性も低下するので、内外面の溶接入熱は1.0〜3.5kJ/mmとした。 In addition, when the welding part of tack welding and inner and outer surface welding overlaps, the one where welding heat input is as low as possible is preferable. Also, the welding heat input varies depending on the plate thickness. However, if the heat input is too small, the penetration becomes insufficient, the number of weldings increases, the work efficiency deteriorates, and if the heat input is too large, the heat-affected zone softens. It is large and the toughness of the welded portion is also reduced, so the welding heat input on the inner and outer surfaces is 1.0 to 3.5 kJ / mm.
シーム溶接後、拡管により真円度を向上させる。真円にするためには塑性域まで変形させる必要がある。本発明の超高強度鋼管の場合は、拡管後円周と拡管前円周の差を拡管前円周で除した値を百分率で表した拡管率が、0.7%以上であることが好ましい。一方、拡管率が2%を超えると、母材、溶接部とも塑性変形により靭性が劣化することがある。したがって、拡管率は0.7〜2%の範囲とすることが好ましい。 After seam welding, roundness is improved by pipe expansion. In order to make a perfect circle, it is necessary to deform to the plastic region. In the case of the ultra-high-strength steel pipe of the present invention, it is preferable that the pipe expansion rate expressed as a percentage obtained by dividing the difference between the circumference after pipe expansion and the circumference before pipe expansion by the circumference before pipe expansion is 0.7% or more. . On the other hand, if the expansion ratio exceeds 2%, the toughness may deteriorate due to plastic deformation of both the base metal and the welded part. Therefore, it is preferable that the tube expansion rate is in the range of 0.7 to 2%.
表1の化学成分からなる鋼を溶製して鋳造し、厚みが240mmの鋼塊とした。これらの鋼塊を1150℃に加熱し、900℃以上で59〜80mm厚さまで再結晶温度域で熱間圧延し、そのまま14〜20mm厚さまで未再結晶域の熱間圧延を600〜880℃で行った。なお、900℃以上は再結晶温度域であり、880℃以下は未再結晶温度域である。熱間圧延後、600〜350℃の範囲を平均冷速で1℃/s以上10℃/s以下で350℃以下まで水冷した。得られた鋼板を筒状にプレス成形し、仮付け溶接を行った後、溶接入熱を2.4kJ/mmとして内外面をサブマージドアーク溶接し、拡管して、36インチ(913mm径)、16mm厚の鋼管とした。表4、表5(表4のつづき)にはこのときの製造条件、母材の特性、試験結果等を示しておく。 Steel made of the chemical components shown in Table 1 was melted and cast into a steel ingot having a thickness of 240 mm. These ingots are heated to 1150 ° C., hot rolled at 900 ° C. or higher to a thickness of 59 to 80 mm in a recrystallization temperature range, and hot rolled in an unrecrystallized zone to a thickness of 14 to 20 mm as it is at 600 to 880 ° C. went. In addition, 900 degreeC or more is a recrystallization temperature range, and 880 degrees C or less is a non-recrystallization temperature range. After hot rolling, the range of 600 to 350 ° C. was water-cooled at an average cooling rate of 1 ° C./s to 10 ° C./s to 350 ° C. or less. The obtained steel plate was press-formed into a cylindrical shape, and after tack welding, the inner and outer surfaces were submerged arc welded with a welding heat input of 2.4 kJ / mm, expanded, and 36 inches (913 mm diameter), The steel pipe was 16 mm thick. Tables 4 and 5 (continued in Table 4) show the manufacturing conditions, characteristics of the base material, test results, and the like at this time.
得られた鋼管の母材から試験片を採取し、JIS Z 2241に準拠して引張試験を行った。また、円周方向を長手としたDWTT試験片を採取し、板厚方向にプレスノッチを導入して、更に3点曲げで延性亀裂を導入し、プリクラックDWTT試験を0℃で実施した。 Test pieces were collected from the obtained steel pipe base material and subjected to a tensile test in accordance with JIS Z 2241. In addition, a DWTT test piece having a longitudinal direction in the circumferential direction was collected, a press notch was introduced in the thickness direction, a ductile crack was further introduced by three-point bending, and a precrack DWTT test was performed at 0 ° C.
鋼1〜8は本発明の例を示す。表4、表5(表4のつづき)から明らかなように、これらの鋼は何れも母材の10℃でのプリクラックDWTTエネルギーが3000J以上である。しかもこれらの鋼管は部分ガスバースト試験で亀裂が停止している。すなわち、不安定延性破壊特性が優れている。それに対し、鋼9〜19は本発明方法から逸脱した比較例を示す。すなわち、鋼9〜13および15は母材の化学成分を逸脱しているかあるいは熱間圧延条件を逸脱しているかによってフェライト/マルテンサイトの2相混合組織になっていない。その結果、D方向の(100)の集積度が3を超えるために母材のプリクラックDWTTエネルギーが3000J未満である。しかも、これらの特性が悪いために部分ガスバースト試験も貫通し、不安定延性破壊特性も劣っている。鋼14は母材のCが逸脱しているので、母材の硬度が満足していない。また、鋼16、17、19は溶接金属の化学成分あるいは入熱条件が逸脱しているために溶接金属強度を満たしていないかあるいは溶接金属靭性を満足していない。鋼18はサブマージドアーク溶接の入熱が低いために溶接欠陥が多発した。 Steels 1-8 show examples of the present invention. As is clear from Tables 4 and 5 (continuation of Table 4), the pre-crack DWTT energy at 10 ° C. of the base material of each of these steels is 3000 J or more. Moreover, cracks of these steel pipes have stopped in the partial gas burst test. That is, the unstable ductile fracture characteristics are excellent. On the other hand, steel 9-19 shows the comparative example which deviated from this invention method. That is, steels 9 to 13 and 15 do not have a ferrite / martensite two-phase mixed structure depending on whether they deviate from the chemical composition of the base metal or the hot rolling conditions. As a result, since the degree of integration of (100) in the D direction exceeds 3, the pre-crack DWTT energy of the base material is less than 3000 J. Moreover, since these characteristics are poor, the partial gas burst test is also penetrated, and the unstable ductile fracture characteristics are also inferior. Since the base material C deviates from the steel 14, the hardness of the base material is not satisfied. Steels 16, 17, and 19 do not satisfy the weld metal strength or do not satisfy the weld metal toughness because the chemical components or heat input conditions of the weld metal deviate. Steel 18 had many welding defects due to the low heat input of submerged arc welding.
表2、表3(表2のつづき)には参考に本発明および比較例の溶接金属および溶接ワイヤーの成分を示した。 Tables 2 and 3 (continued in Table 2) show the components of the weld metal and the weld wire of the present invention and the comparative example for reference.
Claims (13)
C :0.03〜0.10%、
Si:0.01〜0.6%、
Mn:1.5〜2.5%、
P :0.015%以下、
S :0.003%以下、
Ni:0.1〜2.0%、
Mo:0.15〜0.60%、
Nb:0.001〜0.10%、
Ti:0.005〜0.030%、
Al:0.06%以下
を含有し、更に、
B :0.0001〜0.005%、
N :0.0001〜0.006%、
V :0.001〜0.10%、
Cu:0.01〜1.0%、
Cr:0.01〜0.8%、
Zr:0.0001〜0.005%、
Ta:0.0001〜0.005%、
Ca:0.0001〜0.01%、
REM:0.0001〜0.01%、
Mg:0.0001〜0.006%
の1種又は2種以上を含有し、残部が鉄及び不可避的不純物からなることを特徴とする請求項3または4に記載の高速延性破壊特性に優れた超高強度鋼板。 % By mass
C: 0.03-0.10%,
Si: 0.01 to 0.6%,
Mn: 1.5 to 2.5%
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 2.0%,
Mo: 0.15-0.60%,
Nb: 0.001 to 0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less, further,
B: 0.0001 to 0.005%,
N: 0.0001 to 0.006%,
V: 0.001 to 0.10%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 0.8%
Zr: 0.0001 to 0.005%,
Ta: 0.0001 to 0.005%,
Ca: 0.0001 to 0.01%,
REM: 0.0001-0.01%
Mg: 0.0001 to 0.006%
The ultra-high strength steel sheet having excellent high-speed ductile fracture characteristics according to claim 3 or 4, wherein one or more of the above are contained, and the balance consists of iron and inevitable impurities.
C :0.04〜0.14%、
Si:0.05〜0.4%、
Mn:1.2〜2.2%、
P :0.01%以下、
S :0.010%以下、
Ni:1.3〜3.2%、
Cr+Mo+V:1.0〜2.5%、
Ti:0.003〜0.050%、
Al:0.02%以下、
B:0.005%以下、
O:0.01〜0.03%
を含有し、残部が鉄及び不可避的不純物からなることを特徴とする請求項9に記載の高速延性破壊特性に優れた超高強度鋼管。 The composition of the weld metal is mass%,
C: 0.04 to 0.14%,
Si: 0.05-0.4%
Mn: 1.2-2.2%,
P: 0.01% or less,
S: 0.010% or less,
Ni: 1.3-3.2%
Cr + Mo + V: 1.0 to 2.5%,
Ti: 0.003 to 0.050%,
Al: 0.02% or less,
B: 0.005% or less,
O: 0.01 to 0.03%
The ultrahigh strength steel pipe excellent in high-speed ductile fracture characteristics according to claim 9, wherein the balance is made of iron and inevitable impurities.
C :0.01〜0.12%、
Si:0.3%以下、
Mn:1.2〜2.4%、
Ni:4.0〜8.5%、
Cr+Mo+V:3.0〜5.0%、
Ti:0.005〜0.15%、
Al:0.02%以下
を含有し、残部が鉄及び不可避的不純物からなる溶接ワイヤーを用いてサブマージドアーク溶接することを特徴とする請求項11または12に記載の高速延性破壊特性に優れた超高強度鋼管の製造方法。 % By mass
C: 0.01 to 0.12%,
Si: 0.3% or less,
Mn: 1.2-2.4%
Ni: 4.0 to 8.5%,
Cr + Mo + V: 3.0-5.0%,
Ti: 0.005 to 0.15%,
It is excellent in the high-speed ductile fracture property according to claim 11 or 12, characterized in that submerged arc welding is performed using a welding wire containing Al: 0.02% or less, the balance being iron and inevitable impurities. Manufacturing method of ultra high strength steel pipe.
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