WO2011043287A1 - Steel for linepipe having good strength and malleability, and method for producing the same - Google Patents

Steel for linepipe having good strength and malleability, and method for producing the same Download PDF

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WO2011043287A1
WO2011043287A1 PCT/JP2010/067351 JP2010067351W WO2011043287A1 WO 2011043287 A1 WO2011043287 A1 WO 2011043287A1 JP 2010067351 W JP2010067351 W JP 2010067351W WO 2011043287 A1 WO2011043287 A1 WO 2011043287A1
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steel
ductility
less
good strength
strength
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石川 肇
植森 龍治
渡部 義之
侭田 伸彦
潔 海老原
児島 明彦
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新日本製鐵株式会社
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Priority to BR112012007636-1A priority Critical patent/BR112012007636B1/en
Priority to JP2011501826A priority patent/JP4824142B2/en
Publication of WO2011043287A1 publication Critical patent/WO2011043287A1/en

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    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00Heat treatment of ferrous alloys
    • C21D6/005Heat treatment of ferrous alloys containing Mn
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/10Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C21D8/105Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies of ferrous alloys
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/001Ferrous alloys, e.g. steel alloys containing N
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/002Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/02Ferrous alloys, e.g. steel alloys containing silicon
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/04Ferrous alloys, e.g. steel alloys containing manganese
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/06Ferrous alloys, e.g. steel alloys containing aluminium
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/12Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/14Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00Microstructure comprising significant phases
    • C21D2211/005Ferrite
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00Microstructure comprising significant phases
    • C21D2211/009Pearlite

Abstract

Provided is steel for linepipe having good strength and malleability. The steel comprises C: 0.07% to 0.15%, Si: 0.05% to 0.60%, Mn: 0.80% to 1.80%, P: 0.010% or less, S: 0.007% or less, V: 0.05% to 0.12%, Nb: 0.005% to 0.070%, Al: 0.005% to 0.08%, Ti: 0.005% to 0.030%, Ca: 0.0005% to 0.0035%, N: 0.0020% to 0.0060%, O: 0.0030% or less, and iron and inevitable impurities as the balance; the structure is a two-phase structure of ferrite and pearlite; the surface area percentage of island martensite is less than 1.5%; sheet thickness is 18 mm or greater; and yield strength is 450 MPa or greater.

Description

強度、延性の良好なラインパイプ用鋼およびその製造方法Steel for line pipe with good strength and ductility and method for producing the same
 本発明は溶接構造用鋼としての十分な強度を有し、かつ延性特性に優れるとともに低温靭性に優れた高靭性、高強度、高延性鋼およびその製造方法に関するものである。特に、寒冷地で低温靭性が要求される伸びに優れた強度、延性を有するラインパイプ用鋼およびその製造方法に関するものである。
 本願は、2009年10月5日に、日本に出願された特願2009-231799号に基づき優先権を主張し、その内容をここに援用する。
The present invention relates to a high toughness, high strength, high ductility steel having a sufficient strength as a welded structural steel and having excellent ductility characteristics and excellent low temperature toughness, and a method for producing the same. In particular, the present invention relates to a steel for line pipes having strength and ductility excellent in elongation that requires low temperature toughness in a cold region and a method for producing the same.
This application claims priority based on Japanese Patent Application No. 2009-231799 filed in Japan on October 5, 2009, the contents of which are incorporated herein by reference.
 近年、安全性の向上や輸送ガスの高圧化等による操業効率の向上、使用鋼材量の削減によるコストの低減のため、ラインパイプ用鋼に対して高強度化が求められている。また、該鋼材の使用される地域は、寒冷地などの自然環境が苛酷な地域へと拡大しつつあり、厳しい靭性特性が必要とされている。特に、地震多発地域などに使用される構造物用鋼では従来の要求特性に加えて塑性変形能や耐延性破壊特性などが求められている。 In recent years, there has been a demand for higher strength steel for line pipes in order to improve safety, improve operational efficiency by increasing the pressure of transport gas, and reduce costs by reducing the amount of steel used. In addition, the area where the steel material is used is expanding to an area where the natural environment such as a cold region is severe, and severe toughness characteristics are required. In particular, structural steels used in earthquake-prone areas are required to have plastic deformability and ductile fracture characteristics in addition to the conventional required characteristics.
 たとえば、特許文献1では、延性破壊を抑制するために高い一様伸びを得ることを目的とした鋼を提示している。ここでは、焼入、二相処理、焼戻処理(QLT処理)によりフェライト内に適量の硬化相を導入した混合組織を混在させ高延性を図っている。また、特許文献2では、鋼成分と焼入れ硬化性(Di)の最適化と加速冷却により高延性を図っている。 For example, Patent Document 1 presents steel aimed at obtaining high uniform elongation in order to suppress ductile fracture. Here, a mixed structure in which an appropriate amount of a hardened phase is introduced into ferrite is mixed by quenching, two-phase treatment, and tempering treatment (QLT treatment) to achieve high ductility. Moreover, in patent document 2, high ductility is aimed at by optimization and accelerated cooling of a steel component and quenching hardenability (Di).
 特許文献3では、耐HIC特性の優れた鋼板を提示している。16mm以下でX56以下の鋼板を加速冷却せずにフェライト-パーライト-ベイナイトからなる混合組織を製造し耐HIC特性を確保している。
 特許文献4では、570MPa以上の鋼板を提示している。加速冷却後、誘導加熱により島状マルテンサイト(MA、Martensite-Austenite-Constituent)の生成を抑え、表層の硬さを抑えている。板厚方向の硬さのバラツキを抑えることにより高強度高靭性をはかっている。
Patent Document 3 presents a steel sheet having excellent HIC resistance. A mixed structure composed of ferrite-pearlite-bainite is manufactured without accelerated cooling of a steel plate of 16 mm or less and X56 or less to ensure HIC resistance.
Patent Document 4 presents a steel plate of 570 MPa or more. After accelerated cooling, the formation of island martensite (MA) is suppressed by induction heating, and the hardness of the surface layer is suppressed. High strength and toughness are achieved by suppressing variations in hardness in the thickness direction.
 一般に、高強度鋼では強度を得るために炭素当量や焼入れ指数を上昇させることが必要とされている。しかし、単純に炭素当量を上昇させた場合、延性や靭性の低下を招く。一方、大径ラインパイプ用鋼ではUOE、JCOEなど造管後の延性を管理するために板内での強度や延性などのバラツキの低減が要求されており、ミクロ組織での著しい硬化組織の生成の抑制が必要である。 Generally, in order to obtain strength in high-strength steel, it is necessary to increase the carbon equivalent and quenching index. However, when the carbon equivalent is simply increased, ductility and toughness are reduced. On the other hand, steel for large-diameter line pipes is required to reduce variations in strength and ductility within the plate in order to manage ductility after pipe making, such as UOE and JCOE. It is necessary to suppress this.
特開2003-253331号公報JP 2003-253331 A 特開2003-288512号公報JP 2003-288512 A 特開2001-158936号公報JP 2001-158936 A 特開2008-121036号公報JP 2008-121036 A
 本発明者は従来の技術に以下のような問題があることを見出した。
 大径ラインパイプ用鋼ではUOE、JCOEなど造管後の延性を管理するために板内での強度や延性などの特性のバラツキの低減が要求されている。そのため、たとえばQLT(Quenching-Lamellarizing-Tempering)処理による均一な組織の形成により鋼特性の板内バラツキを小さくする技術が採用されている。しかし、QLT処理は少なくとも高温で3回以上の熱処理をするため高コストとなる。また、二相域熱処理に相当する加速冷却によって高強度、高延性をはかることは可能であるが、加速冷却において板内の冷却を均一にすることは極めて困難である。
The present inventor has found that the prior art has the following problems.
Steel for large diameter line pipes is required to reduce variations in properties such as strength and ductility in the plate in order to manage ductility after pipe making such as UOE and JCOE. For this reason, for example, a technique for reducing the in-plate variation in steel characteristics by forming a uniform structure by QLT (Quenching-Lamellarizing-Tempering) processing is adopted. However, the QLT process is expensive because it is heat-treated at least three times at a high temperature. Further, it is possible to achieve high strength and high ductility by accelerated cooling corresponding to the two-phase region heat treatment, but it is extremely difficult to make the cooling in the plate uniform in accelerated cooling.
 上記特許文献3のように耐HIC特性を確保するためにはNACEで規定されているようにビッカーズ硬さで248Hv以下と、鋼全体としての硬さの低減が求められる。このため、高強度化は困難である。また、このような鋼を加速冷却で製造すると硬化組織であるベイナイトやMAの生成は避けられない。この場合、伸び(特に局部伸び)の低下は避けられない。 In order to ensure the HIC resistance as described in Patent Document 3, it is required to reduce the hardness of the steel as a whole by Vickers hardness of 248 Hv or less as specified by NACE. For this reason, it is difficult to increase the strength. Moreover, when such steel is manufactured by accelerated cooling, the production | generation of bainite and MA which are hardened structures is inevitable. In this case, a decrease in elongation (particularly local elongation) is inevitable.
 上記特許文献4でのように加速冷却後、誘導加熱によって表層の硬さを低減させることによって、局部伸びは向上できる。しかしながら、加速冷却により製造され、ベイナイトとフェライトの混合組織と称されている組織には、該公報の図2に示されるように、バンド組織が形成されず、一般的なフェライトは認められない。このような組織では局部伸びはよくなるが、一様伸びは著しく低下させ、かえって全伸びは低下する。また、製造ラインに誘導加熱設備を設置することはコストがかかるとともに誘導加熱による熱処理は電磁波での表皮の深さの関係から制御が著しく難しい。 The local elongation can be improved by reducing the hardness of the surface layer by induction heating after accelerated cooling as in Patent Document 4 above. However, in a structure manufactured by accelerated cooling and called a mixed structure of bainite and ferrite, as shown in FIG. 2 of the publication, no band structure is formed, and general ferrite is not recognized. In such a structure, the local elongation is improved, but the uniform elongation is significantly reduced, and the total elongation is reduced. In addition, it is expensive to install induction heating equipment in the production line, and heat treatment by induction heating is extremely difficult to control because of the depth of the skin due to electromagnetic waves.
 そこで、本発明ではラインパイプ用鋼において靭性、延性特性の良好な廉価な高強度厚鋼板とその製造方法、およびその品質管理方法を提供するものである。 Therefore, the present invention provides an inexpensive high-strength thick steel plate having good toughness and ductility in line pipe steel, a manufacturing method thereof, and a quality control method thereof.
 本発明者らは、延性におよぼすMAなどの硬化組織の影響やMAの生成を助長する偏析について調査した結果、強度、延性バランスの観点から、鋼の組織をフェライトとパーライトの二相組織とし、かつ、MAの生成を抑制することにより、局部伸びの低下を防止することを見出した。 As a result of investigating the influence of hardened structure such as MA on the ductility and segregation that promotes the formation of MA, the present inventors have made the steel structure a two-phase structure of ferrite and pearlite from the viewpoint of strength and ductility balance. And it discovered that the fall of local elongation was prevented by suppressing the production | generation of MA.
 本発明の要旨は、次のとおりである。
(1)本発明の一態様にかかる強度および延性の良好なラインパイプ用鋼は、質量%で、C:0.07~0.15%、Si:0.05~0.60%、Mn:0.80~1.80%、P:0.010以下、S:0.007%以下、V:0.05~0.12%、Nb: 0.005~0.070%、Al:0.005~0.08%、Ti:0.005~0.030%、Ca:0.0005~0.0035%、N:0.0020~0.0060%、O:0.0030%以下を含有し、残部が鉄および不可避的不純物からなり、組織がフェライトとパーライトの二相組織であり、島状マルテンサイトの面積分率が1.5%未満であり、板厚が18mm以上であり、降伏強度が450MPa以上である。
(2)上記(1)のラインパイプ用鋼は、更に、質量%で、Cu:0.05~0.70%、Ni:0.05~0.70%、Cr:0.80%以下、Mo:0.30%以下、B:0.0003~0.0030%、Mg:0.0003~0.0030%、REM:0.0005~0.0050%の一種または二種以上を含有してもよい。
(3)上記(1)または(2)のラインパイプ用鋼において、光学顕微鏡下でベイナイト組織が検出されなくてもよい。
(4)上記(1)または(2)のラインパイプ用鋼において、Mnの偏析度が1.7以下であってもよい。
(5)上記(4)のラインパイプ用鋼において、Si,Pの偏析度がそれぞれ1.5以下、8.0以下であってもよい。
(6)本発明の別の一態様にかかる強度および延性の良好なラインパイプ用鋼の製造方法は、上記(1)または(2)のラインパイプ用鋼の化学成分を持つスラブを1250℃以下の温度に加熱後、850℃以上の温度域において累積圧下率で40%以上の熱間圧延をし、700~800℃の温度領域で熱間圧延を終了させた後、空冷する。
(7)上記(6)の製造方法において、前記空冷後、鋼板に500~300℃の温度領域で焼戻し処理を施してもよい。
(8)本発明の別の一態様にかかる強度および延性の良好なラインパイプ用鋼の品質管理方法は、熱間圧延によって製造された鋼板中の島状マルテンサイト量を測定し、島状マルテンサイトの面積分率が1.5%未満になるように製造条件を制御することによって延性を改善させる。
(9)上記(8)の品質管理方法では、熱間圧延後の冷却方法を空冷とすることによって延性を改善させてもよい。
(10)上記(8)または(9)の品質管理方法では、鋼板段階のMnの偏析度を1.7以下とすることによって延性を改善させてもよい。
(11)上記(8)または(9)の品質管理方法では、鋼板段階のベイナイト組織分率を1%以下にすることによって延性を改善させてもよい。
The gist of the present invention is as follows.
(1) The steel for line pipes with good strength and ductility according to one embodiment of the present invention is C: 0.07 to 0.15%, Si: 0.05 to 0.60%, Mn: 0.80 to 1.80%, P: 0.010 or less, S: 0.007% or less, V: 0.05 to 0.12%, Nb: 0.005 to 0.070%, Al: 0.00. 005 to 0.08%, Ti: 0.005 to 0.030%, Ca: 0.0005 to 0.0035%, N: 0.0020 to 0.0060%, O: 0.0030% or less The balance is composed of iron and inevitable impurities, the structure is a two-phase structure of ferrite and pearlite, the area fraction of island martensite is less than 1.5%, the plate thickness is 18 mm or more, and the yield strength Is 450 MPa or more.
(2) The steel for line pipes of (1) above is further, in mass%, Cu: 0.05 to 0.70%, Ni: 0.05 to 0.70%, Cr: 0.80% or less, One or more of Mo: 0.30% or less, B: 0.0003 to 0.0030%, Mg: 0.0003 to 0.0030%, REM: 0.0005 to 0.0050% are contained. Also good.
(3) In the steel for line pipes according to (1) or (2), the bainite structure may not be detected under an optical microscope.
(4) In the steel for line pipes of (1) or (2) above, the segregation degree of Mn may be 1.7 or less.
(5) In the steel for line pipes of (4) above, the segregation degrees of Si and P may be 1.5 or less and 8.0 or less, respectively.
(6) According to another aspect of the present invention, a method for producing a steel for a line pipe with good strength and ductility is obtained by adding a slab having a chemical component of the steel for a line pipe according to (1) or (2) to 1250 ° C. After heating to a temperature of 850 ° C., hot rolling is performed at a cumulative reduction ratio of 40% or higher in a temperature range of 850 ° C. or higher, and hot rolling is terminated in a temperature range of 700 to 800 ° C., followed by air cooling.
(7) In the manufacturing method of (6), the steel sheet may be tempered in the temperature range of 500 to 300 ° C. after the air cooling.
(8) According to another aspect of the present invention, a method for quality control of steel for line pipes with good strength and ductility measures the amount of island martensite in a steel plate produced by hot rolling, Ductility is improved by controlling the production conditions so that the area fraction is less than 1.5%.
(9) In the quality control method of (8) above, the ductility may be improved by setting the cooling method after hot rolling to air cooling.
(10) In the quality control method of (8) or (9) above, the ductility may be improved by setting the segregation degree of Mn at the steel plate stage to 1.7 or less.
(11) In the quality control method of the above (8) or (9), the ductility may be improved by setting the bainite structure fraction at the steel sheet stage to 1% or less.
 以上述べたように、本発明によれば強度、延性の良好なラインパイプ用鋼が得られるため、産業上極めて有用なものである。 As described above, according to the present invention, steel for line pipes having good strength and ductility can be obtained, which is extremely useful industrially.
MAと、一様伸び(U.El)、局部伸び(L.El)および全伸び(T.El)の関係を示すグラフである。It is a graph which shows the relationship between MA, uniform elongation (U.El), local elongation (L.El), and total elongation (T.El).
 一般に、高強度化には多量の合金添加や加速冷却は有効であるが、焼入れ性の高い組織となるためにかえって延性を劣化させる。また、高強度化した場合、局部的にMAや焼きの入ったベイナイトなどが生成する場合があるが、これらと伸びとの関係は不明確な点が多い。 In general, a large amount of alloy addition and accelerated cooling are effective for increasing the strength, but the ductility deteriorates on the contrary because of a highly hardenable structure. In addition, when the strength is increased, MA or baked bainite may be locally generated, but the relationship between these and elongation is unclear.
 そこで、本発明者らは延性におよぼす組織の影響について詳細な研究を実施し、延性におよぼすMAなどの硬化組織の影響やMAの生成を助長する偏析について調査し、以下のことが必要であることを明らかにした。
 1. 強度、延性バランスの観点からフェライトとパーライトの二相組織とする必要がある。
 2. MAの生成は一様伸びに及ぼす影響は少ないが、MA分率が上がると局部伸びが著しく劣化する。また、約400Hv以上のベイナイトもMAと同様の挙動を示し、その生成は局部伸びを低下させる。
 3. MA生成を抑制するためには、製造プロセスを制御することも重要であるが、偏析の低減も重要である。特に偏析については偏析度が高くなるとバンド状のMAを形成し著しく局部伸びを低下させる。
Therefore, the present inventors have conducted detailed studies on the influence of the structure on the ductility, investigated the influence of the hardened structure such as MA on the ductility and the segregation that promotes the formation of MA, and the following is necessary. It revealed that.
1. From the viewpoint of strength and ductility balance, it is necessary to have a two-phase structure of ferrite and pearlite.
2. The production of MA has little influence on the uniform elongation, but as the MA fraction increases, the local elongation deteriorates significantly. Also, bainite of about 400 Hv or more shows the same behavior as MA, and its generation reduces local elongation.
3. In order to suppress the formation of MA, it is important to control the manufacturing process, but it is also important to reduce segregation. In particular, as for segregation, when the degree of segregation increases, band-like MA is formed and the local elongation is remarkably reduced.
 前述のように一般にラインパイプ用として高強度化をはかった材料の延性は低値となる。たとえば、加速冷却を用いてベイナイト単相組織とした場合、YS(降伏強度)で450MPa級程度の強度確保は容易である。しかし、強度の上昇により延性(一様伸び)は低下する。また、単相組織とした場合、延性に関しては特に局部伸びが著しく低下し、強度・延性バランスの確保は困難である。また、フェライト単相とした場合、高延性化することは可能となるが強度の確保は難しい。また、フェライトと加速冷却によって生ずる硬化したべイナイト組織とがあると局部伸びを低下させる。
 このため、高延性化をはかるためのフェライトと強度を確保するためのパーライトの二相組織が必要となる。
As described above, the ductility of a material which is generally increased in strength for line pipes is low. For example, when a bainite single phase structure is formed by using accelerated cooling, it is easy to secure a strength of about 450 MPa class in terms of YS (yield strength). However, the ductility (uniform elongation) decreases with increasing strength. In the case of a single-phase structure, particularly with respect to ductility, local elongation is remarkably reduced, and it is difficult to ensure a balance between strength and ductility. Further, when a ferrite single phase is used, it is possible to increase ductility, but it is difficult to ensure strength. Further, if there is ferrite and a hardened bainitic structure generated by accelerated cooling, local elongation is lowered.
For this reason, a two-phase structure of ferrite for ensuring high ductility and pearlite for ensuring strength is required.
 また、強度化した場合、パーライトまたはベイナイトの一部にMAが生成することは避けられない。このMAと一様伸び(U.El)、局部伸び(L.El)および全伸び(T.El)の関係を示した一例を図1に示す。この図1に示すにように、MAの増加により、一様伸びはほとんど低下しないが局部伸びが低下する。MAの増加による局部伸びの低下の結果、全伸びが低下することが判る。一様伸びの低下を避けるため、本発明では、MAを1.5%以下に制限している。MAを1.0%以下または0.5%以下に制限することがより好ましい。 Also, when strengthened, it is inevitable that MA is generated in part of pearlite or bainite. An example showing the relationship between this MA and uniform elongation (U. El), local elongation (L. El) and total elongation (T. El) is shown in FIG. As shown in FIG. 1, due to the increase in MA, the uniform elongation hardly decreases but the local elongation decreases. It can be seen that the overall elongation decreases as a result of the decrease in local elongation due to the increase in MA. In order to avoid a decrease in uniform elongation, the MA is limited to 1.5% or less in the present invention. More preferably, MA is limited to 1.0% or less or 0.5% or less.
 よって、本発明では廉価にしつつ、フェライトとパーライトの二相組織の制御とMAおよび偏析の制御により高強度鋼にも関わらず延性の確保をはかったものである。 Therefore, in the present invention, the ductility is ensured despite the high-strength steel by controlling the two-phase structure of ferrite and pearlite and controlling the MA and segregation, while being inexpensive.
 以下に本発明の構成について詳細に説明する。まず、本発明の鋼材の組成限定理由について説明する。なお、本明細書においては、含有率%は、全て質量%を意味する。 Hereinafter, the configuration of the present invention will be described in detail. First, the reasons for limiting the composition of the steel material of the present invention will be described. In the present specification, the content% means mass%.
C:0.07~0.15%
 Cは強度を確保するために必要な元素であり、0.07%以上の添加が必要である。多量の添加は延性や低温靭性の低下を招くおそれがあるために、その上限値を0.15%とする。望ましくは0.12%以下が良い。
C: 0.07 to 0.15%
C is an element necessary for ensuring strength, and it is necessary to add 0.07% or more. Since a large amount of addition may cause a decrease in ductility and low temperature toughness, the upper limit is set to 0.15%. Desirably, 0.12% or less is good.
Si:0.05~0.60%
 Siは脱酸元素として、また固溶強化により鋼の強度を増加させるのに有効な元素であるが、0.05%未満の添加ではそれらの効果が認められない。また、0.60%を超えて添加すると、組織内にMAが多量に生成するため靭性を劣化させる。このため、Siの添加量は0.20~0.60%とした。なお、0.45%以上になるとMA(または残留オーステナイト)が増加し始めるため、望ましくは0.45%未満が良い。
Si: 0.05 to 0.60%
Si is an element effective as a deoxidizing element and increases the strength of the steel by solid solution strengthening. However, when its content is less than 0.05%, these effects are not recognized. Further, if added over 0.60%, a large amount of MA is generated in the structure, so that the toughness is deteriorated. Therefore, the addition amount of Si is set to 0.20 to 0.60%. Note that since MA (or retained austenite) starts to increase at 0.45% or more, it is preferably less than 0.45%.
Mn:0.80~1.80%
 Mnは、鋼の強度を増加するため高強度化には有効な元素である。そのためには、0.80%以上の添加が必要である。しかし、1.80%を超えると、中心偏析やミクロ偏析等の偏析度が増加するためMAの生成を助長するため局部伸びを低下させる。このため、Mnの添加量の適正範囲を0.80~1.80%とした。高強度化のため、Mnの下限を1.0%、1.2%または1.3%とすることが望ましい。
Mn: 0.80 to 1.80%
Mn is an effective element for increasing strength because it increases the strength of steel. For that purpose, addition of 0.80% or more is necessary. However, if it exceeds 1.80%, the degree of segregation such as center segregation and microsegregation increases, so that the formation of MA is promoted and the local elongation is lowered. Therefore, the appropriate range of Mn addition amount is set to 0.80 to 1.80%. In order to increase the strength, it is desirable that the lower limit of Mn is 1.0%, 1.2% or 1.3%.
P:0.010%以下
 Pは、0.010%超となると粒界に偏析して鋼の靱性を著しく劣化させる。また、偏析帯のPの濃度が増加しMAの生成を助長する。このため添加量の上限を0.010%とした。なお、延性や靭性値の低下を抑制する観点からはできるだけ低減することが望ましい。
P: 0.010% or less When P exceeds 0.010%, it segregates at the grain boundaries and significantly deteriorates the toughness of the steel. In addition, the concentration of P in the segregation zone is increased to promote the formation of MA. For this reason, the upper limit of the addition amount was set to 0.010%. In addition, it is desirable to reduce as much as possible from a viewpoint of suppressing the fall of a ductility and a toughness value.
S:0.007%以下
 Sは、MnSを形成して鋼中に存在し、圧延冷却後の組織を微細にする作用を有するが、0.007%を超えると母材および溶接部の靭性を劣化させる。このため、Sは0.007%以下とした。
S: 0.007% or less S is present in steel by forming MnS, and has the effect of refining the structure after rolling and cooling. However, if it exceeds 0.007%, the toughness of the base metal and the welded portion is reduced. Deteriorate. For this reason, S was made into 0.007% or less.
Nb:0.005~0.070%
 本発明では組織制御のため冷却を空冷としたため、強度を確保するためにNbが重要な元素である。また、Nb添加によって、スラブ再加熱時や焼入れ時の加熱オーステナイトが細粒化し、鋼の高強度化がはかれる。そのためには0.005%以上添加する必要がある。しかしながら、過量なNb添加は母材の延性を低下させるため、Nb添加量の上限値を0.070%とした。母材の靭性向上のために、Nbの上限を0.050%または0.35%に制限してもよい。
Nb: 0.005 to 0.070%
In the present invention, since the cooling is air cooling for the structure control, Nb is an important element for securing the strength. Moreover, by adding Nb, the heated austenite at the time of slab reheating or quenching becomes finer and the strength of the steel is increased. Therefore, it is necessary to add 0.005% or more. However, since excessive Nb addition reduces the ductility of the base material, the upper limit of the Nb addition amount is set to 0.070%. In order to improve the toughness of the base material, the upper limit of Nb may be limited to 0.050% or 0.35%.
V:0.05~0.12%
 Vは、Nbとほぼ同様の作用を有するが、Nbに比べてその効果は小さい。Nbと同様の効果は0.05%未満では効果が少ない。しかし、0.12%を超えると延性が劣化する。このため、Vの添加量の適正範囲を0.05~0.12%とした。
V: 0.05 to 0.12%
V has almost the same function as Nb, but its effect is smaller than that of Nb. The effect similar to Nb is less effective at less than 0.05%. However, if it exceeds 0.12%, the ductility deteriorates. For this reason, the appropriate range of the amount of V added is set to 0.05 to 0.12%.
Al:0.005~0.08%
 脱酸の目的で、Alは0.005%以上添加する必要がある。0.005%未満とするとMAの生成は抑制されるが、弱脱酸となり粗大な酸化物の発生確率が高くなるため局部伸びを低下させる。一方、0.08%を超える過度の添加は溶接性を低下させる。この問題は、特にフラックスを使用するSAW(Submerged Arc Welding)等で顕著であり、溶接金属の靭性を劣化させ、HAZ(Heart Affected Zone)靱性も低下する。このため、Alの上限を0.08%とした。溶接性の向上のために、Alの上限を0.06または0.04%に制限することが好ましい。
Al: 0.005 to 0.08%
For the purpose of deoxidation, Al needs to be added in an amount of 0.005% or more. If it is less than 0.005%, the formation of MA is suppressed, but it becomes weak deoxidation and the generation probability of coarse oxides is increased, so the local elongation is lowered. On the other hand, excessive addition exceeding 0.08% reduces weldability. This problem is particularly remarkable in SAW (Submerged Arc Welding) using flux, etc., which deteriorates the toughness of the weld metal and decreases the HAZ (Heart Affected Zone) toughness. For this reason, the upper limit of Al was made 0.08%. In order to improve weldability, it is preferable to limit the upper limit of Al to 0.06 or 0.04%.
Ti:0.005~0.030%
 Tiは、Nと結合して鋼中に高強度、高延性化に有効なTiNを形成させるために、0.005%以上の添加が望まれる。ただし、0.030%を超えてTiを添加すると、TiNを粗大化させ、母材の延性を低下させるおそれがある。このため、Tiは0.005~0.030%の範囲とした。
Ca: 0.0005~0.0035%、
 Caは、0.0005%以上添加した場合、硫化物(MnS)の形態を制御し、シャルピーの吸収エネルギーを増大させ低温靭性を向上させる効果がある。ただし、0.0035%を超えると粗大なCaOやCaSが多量に発生するため鋼の延性や靱性に悪影響を及ぼす。このため、ため、0.0035%をCa量の上限と限定した。
Ti: 0.005 to 0.030%
Addition of 0.005% or more is desirable for Ti to combine with N to form TiN effective for high strength and high ductility in steel. However, if Ti is added in an amount exceeding 0.030%, TiN may be coarsened and the ductility of the base material may be reduced. Therefore, Ti is set in the range of 0.005 to 0.030%.
Ca: 0.0005 to 0.0035%,
When Ca is added in an amount of 0.0005% or more, it has an effect of controlling the form of sulfide (MnS), increasing the absorbed energy of Charpy and improving the low temperature toughness. However, if it exceeds 0.0035%, a large amount of coarse CaO and CaS is generated, which adversely affects the ductility and toughness of the steel. For this reason, 0.0035% was limited to the upper limit of the Ca content.
N:0.0020~0.0060%
 Nは、Tiと結合して鋼中に高強度、高延性化に有効なTiNを形成させるために、0.0020%以上の添加が必要である。ただし、Nは固溶強化元素としても非常に大きな効果があるため、多量に添加すると延性を劣化するおそれが考えられる。そのため、延性に大きな影響を与えずTiNの効果を最大限に得られるように、Nの上限を0.0060%とした。
N: 0.0020 to 0.0060%
N needs to be added in an amount of 0.0020% or more in order to form TiN effective in increasing the strength and ductility in steel by bonding with Ti. However, since N has a very large effect as a solid solution strengthening element, adding a large amount may cause deterioration of ductility. Therefore, the upper limit of N is set to 0.0060% so that the effect of TiN can be maximized without greatly affecting the ductility.
O: 0.0030%以下
 高強度鋼の場合、Oが0.0030%を超えると鋼の清浄度、靱性劣化を招く。このため上限値を0.0030%とした。
O: 0.0030% or less In the case of high-strength steel, if O exceeds 0.0030%, the cleanliness and toughness of the steel are deteriorated. Therefore, the upper limit value is set to 0.0030%.
 本発明での基本成分は以上の通りであり、十分に目標値を達成できるが、さらに特性を高めるために、必要に応じて以下の元素のうち一種または二種以上を選択元素として添加してもよい。 The basic components in the present invention are as described above and can sufficiently achieve the target value, but in order to further improve the characteristics, one or more of the following elements are added as selective elements as necessary. Also good.
Cu:0.05~0.70%
 Cuは高強度化をはかるために有利な元素である。Cuによる析出効果を確保するためには0.05%以上の添加が望ましい。しかし過剰な添加は母材の硬さを上昇させ延性を低下させるためその上限を0.70%とした。延性向上のため、Cuを0.50%以下、0.30%以下または0.20%以下に制限することが好ましい。
Cu: 0.05 to 0.70%
Cu is an element advantageous for increasing the strength. In order to ensure the precipitation effect by Cu, addition of 0.05% or more is desirable. However, excessive addition increases the hardness of the base material and decreases the ductility, so the upper limit was made 0.70%. In order to improve ductility, it is preferable to limit Cu to 0.50% or less, 0.30% or less, or 0.20% or less.
Ni:0.05~0.70%
 Niは溶接性等に悪影響をおよぼすことなく、強度、靭性を向上させるほか、Cu割れの防止にも効果がある。これらの効果を得るためには0.05%以上の添加が必要である。しかし、Niは高価であるため0.70%以上とすると廉価に鋼を製造できなくなるため添加量を0.70%以下に制限する。延性向上のためにはNi添加は少ない方が望ましく、0.50%以下、0.30%以下または0.20%以下に制限してもよい。
Ni: 0.05-0.70%
Ni improves strength and toughness without adversely affecting weldability and the like, and is also effective in preventing Cu cracking. In order to obtain these effects, addition of 0.05% or more is necessary. However, since Ni is expensive, if it is made 0.70% or more, it becomes impossible to manufacture steel at a low cost, so the addition amount is limited to 0.70% or less. In order to improve ductility, it is desirable to add less Ni, and the content may be limited to 0.50% or less, 0.30% or less, or 0.20% or less.
Cr:0.80%以下
 Crは母材の強度を高める元素である。しかし、0.80%を超えると母材の硬さを上昇させ延性を劣化させる。そのため上限値を0.80%とした。延性向上のため、Crを0.50%以下、0.30%以下または0.10%以下に制限することが好ましい。
Cr: 0.80% or less Cr is an element that increases the strength of the base material. However, if it exceeds 0.80%, the hardness of the base material is increased and the ductility is deteriorated. Therefore, the upper limit is set to 0.80%. In order to improve ductility, it is preferable to limit Cr to 0.50% or less, 0.30% or less, or 0.10% or less.
Mo:0.30%以下
 CrもMoと同様、母材の強度を高める元素である。しかし、0.30%を超えると母材の硬さを上昇させ延性を劣化させる。そのため上限値を0.30%とした。延性向上のため、Moを0.20%以下または0.10%以下に制限することが好ましい。
Mo: 0.30% or less Cr, like Mo, is an element that increases the strength of the base material. However, if it exceeds 0.30%, the hardness of the base material is increased and the ductility is deteriorated. Therefore, the upper limit is set to 0.30%. In order to improve ductility, it is preferable to limit Mo to 0.20% or less or 0.10% or less.
B:0.0003~0.0030%
 Bは鋼中に固溶して焼入れ性を高め強度を上昇させる元素である。この効果を得るためには0.0003%以上の添加が必要である。しかし、Bを過多に添加すると母材靭性を低下させる。このためその上限値を0.0030%とした。
B: 0.0003 to 0.0030%
B is an element that dissolves in steel to increase the hardenability and increase the strength. In order to obtain this effect, addition of 0.0003% or more is necessary. However, if B is added excessively, the base material toughness is lowered. For this reason, the upper limit was made 0.0030%.
Mg:0.0003~0.0030%
 Mgは、オーステナイト粒の成長をも抑制し、細粒に保つ作用があり、靭性を向上させる。この効果を享受するためには、少なくとも0.0003%以上の添加が必須であり、この量を下限とした。一方、必要以上に添加量が増えても添加量に対する効果代が小さくなるばかりでなく、Mgは製鋼歩留まりが必ずしも高くないため、経済性も失することになる。これらを鑑み、本願発明においては上限を0.0030%に限定した。
Mg: 0.0003 to 0.0030%
Mg also suppresses the growth of austenite grains, has the effect of maintaining fine grains, and improves toughness. In order to enjoy this effect, addition of at least 0.0003% or more is essential, and this amount is set as the lower limit. On the other hand, if the amount added is increased more than necessary, not only does the effect margin for the amount added become small, but Mg does not necessarily have a high steelmaking yield, so the economy is lost. In view of these, the upper limit is limited to 0.0030% in the present invention.
REM:0.0005~0.0050%
 REMもMgと同様、オーステナイト粒の成長をも抑制し、細粒に保つ作用があり、靭性を向上させる。この効果を享受するためには、少なくとも0.0005%以上の添加が必須であり、この量を下限とした。一方、必要以上に添加量が増えても添加量に対する効果代が小さくなるばかりでなく、Mgは製鋼歩留まりが必ずしも高くないため、経済性も失することになる。これらを鑑み、本願発明においては上限を0.0050%に限定した。
REM: 0.0005 to 0.0050%
REM, like Mg, also suppresses the growth of austenite grains, keeps them fine, and improves toughness. In order to enjoy this effect, at least 0.0005% or more must be added, and this amount is set as the lower limit. On the other hand, if the amount added is increased more than necessary, not only does the effect margin for the amount added become small, but Mg does not necessarily have a high steelmaking yield, so the economy is lost. In view of these, the upper limit is limited to 0.0050% in the present invention.
 本発明では主としてラインパイプ溶接用鋼材として高強度、高延性のUOEやJCOE鋼管の製造を可能とする。本発明鋼ではラインパイプに要求される強度、靭性、延性の複合特性を主としてフェライト(フェライトバンド組織を有する)とパーライトの二相組織により確保したことに一つの特徴がある。本発明におけるフェライト組織は、組織内に熱間圧延方向に沿った、いわゆるフェライトバンド組織を有する。 In the present invention, it is possible to manufacture UOE and JCOE steel pipes having high strength and high ductility mainly as steel materials for line pipe welding. One feature of the steel according to the present invention is that the composite properties of strength, toughness and ductility required for the line pipe are ensured mainly by a two-phase structure of ferrite (having a ferrite band structure) and pearlite. The ferrite structure in the present invention has a so-called ferrite band structure along the hot rolling direction in the structure.
 更に、この時、MA分率が1.5%以上となると引張試験時にMA近傍に多量のボイドが発生し、塑性流動によりせん断的な破壊を助長し局部伸びを著しく低下させる。局部伸びを劣化させる偏析起因のMAを抑制し、1.5%以下のMAにするためには急速な冷却をしないことが重要である。より詳細には、ビッカーズ硬さ400~700HvのMAがボイドの発生を特に頻繁に引き起こし、局部伸びを著しく劣化させる要因となる。このため、400~700HvのMAを抑制すれば、局部伸びの劣化を避けることができる。MA分率が上記の範囲である限り、本発明では鋼板中の各元素の偏析度については必ずしも規制しない。しかし、望ましくはMnの鋼板の偏析度を1.7以下にするとよい。さらに望ましくはSi,Pの鋼板の偏析度をそれぞれ1.5、8.0以下にするとよい。
 鋼板のMn偏析度が1.7超、またはSi偏析度が1.5超、またはP偏析度が8.0超になるとMAの生成が顕著となる。本発明では圧延後の鋼板の元素偏析度のみを規定するが、板とした時にMnの偏析度が1.7以下を確保するためには、スラブ段階のMn偏析度は1.1以下にする必要がある。また、本発明では最終的な鋼板のMn偏析を制御するためスラブの製造方法については限定しない。ただし、MAの分布を制御するためには中心偏析などのマクロ偏析だけではなくミクロ偏析も低減する必要がある。
Further, at this time, if the MA fraction is 1.5% or more, a large amount of voids are generated in the vicinity of the MA during the tensile test, and the shear elongation is promoted by the plastic flow, and the local elongation is remarkably reduced. In order to suppress the segregation-induced MA that deteriorates the local elongation and to make the MA less than 1.5%, it is important not to cool rapidly. More specifically, MA having a Vickers hardness of 400 to 700 Hv causes the occurrence of voids particularly frequently and causes a significant deterioration in local elongation. For this reason, if the MA of 400 to 700 Hv is suppressed, deterioration of local elongation can be avoided. As long as the MA fraction is in the above range, the segregation degree of each element in the steel sheet is not necessarily regulated in the present invention. However, the segregation degree of the Mn steel plate is preferably 1.7 or less. More preferably, the segregation degrees of the Si and P steel sheets are 1.5 and 8.0 or less, respectively.
When the Mn segregation degree of the steel sheet exceeds 1.7, the Si segregation degree exceeds 1.5, or the P segregation degree exceeds 8.0, the formation of MA becomes significant. In the present invention, only the element segregation degree of the steel sheet after rolling is specified, but in order to ensure that the Mn segregation degree is 1.7 or less when the sheet is formed, the Mn segregation degree in the slab stage is 1.1 or less. There is a need. Moreover, in this invention, in order to control Mn segregation of the final steel plate, the manufacturing method of a slab is not limited. However, in order to control the MA distribution, it is necessary to reduce not only macro segregation such as center segregation but also micro segregation.
 本発明において、偏析度とは、板厚方向1/2の1mmの領域における成分分析を行い、これらの位置におけるMn、SiまたはPのピークの濃度を各元素の平均濃度で割ったものをいう。成分分析は、例えば、EPMA(Electron Probe Micro Analyzer)またはCMA(Computer-aided Micro Analyzer)を使用することができる。 In the present invention, the degree of segregation means component analysis in a 1 mm 2 region in the plate thickness direction 1/2, and the peak concentration of Mn, Si or P at these positions divided by the average concentration of each element. Say. For component analysis, for example, EPMA (Electron Probe Micro Analyzer) or CMA (Computer-aided Micro Analyzer) can be used.
 次に、本発明鋼材の製造条件限定の理由について説明する。 Next, the reason for limiting the production conditions of the steel of the present invention will be described.
 スラブのMn偏析度を1.1以上とするのが好ましい。このようなスラブから鋼板を製造すれば、鋼板のMnの偏析度を確実に1.7以下にできる。前述のように1.7超となるとMAの生成が顕著となりやすい。なお、本発明ではスラブに対してMnの偏析度のみを規定する。Si,Pの偏析も重要ではあるが、Mnと比較すると製造プロセスの影響が大きいため規定しない。製造プロセス上許容される場合は、Si,Pの偏析は1.5、8.0以下が望ましい。スラブのMn偏析度を制御する方法として、例えば、軽圧下(ソフトリダクション)、連続鋳造時の電磁攪拌、スラブへの高温熱処理による偏析元素の拡散熱処理、など、広く知られている方法を用いることができる。本発明に用いたスラブの製造工程では、軽圧下を行った。 The Mn segregation degree of the slab is preferably 1.1 or more. If a steel plate is manufactured from such a slab, the segregation degree of Mn of a steel plate can be reliably made 1.7 or less. As described above, when it exceeds 1.7, the generation of MA tends to be remarkable. In the present invention, only the segregation degree of Mn is defined for the slab. Although segregation of Si and P is important, it is not specified because the influence of the manufacturing process is larger than that of Mn. If allowed by the manufacturing process, the segregation of Si and P is preferably 1.5 or 8.0 or less. As a method for controlling the Mn segregation degree of the slab, for example, a widely known method such as light reduction (soft reduction), electromagnetic stirring during continuous casting, diffusion heat treatment of the segregation element by high-temperature heat treatment to the slab, and the like should be used. Can do. In the manufacturing process of the slab used in the present invention, light reduction was performed.
 加熱温度が1250℃を超えると、結晶粒径の粗大化が著しく、また、加熱によるスケールが鋼表面に多量に発生し表面の品質が著しく低下する。このため再加熱温度の上限を1250℃とした。
 850℃以上の温度域において累積圧下率で40%以上の熱間圧延を行う必要がある。この温度域における圧下量の増加は、圧延中のオーステナイト粒の微細化に寄与し、結果としてフェライト粒を微細化し機械的性質を向上させる効果がある。このような効果を得るためには、850℃以上の温度域において累積圧下率が40%以上必要である。
When the heating temperature exceeds 1250 ° C., the crystal grain size becomes very coarse, and a large amount of scale is generated on the steel surface, and the quality of the surface is significantly lowered. For this reason, the upper limit of reheating temperature was 1250 degreeC.
It is necessary to perform hot rolling with a cumulative rolling reduction of 40% or more in a temperature range of 850 ° C. or higher. The increase in the amount of reduction in this temperature range contributes to the refinement of the austenite grains during rolling, and as a result, has the effect of refining the ferrite grains and improving the mechanical properties. In order to obtain such an effect, the cumulative rolling reduction needs to be 40% or more in the temperature range of 850 ° C. or higher.
 未再結晶温度域において累積圧下率で40%以上の熱間圧延を行う必要がある。未再結晶温度域における圧下量の増加は、圧延中のオーステナイト粒の微細化に寄与し、結果としてフェライト粒を微細化し機械的性質を向上させる効果がある。このような効果を得るためには、未再結晶域での累積圧下率が40%以上必要である。このため、未再結晶域での累積圧下量を40%以上に限定した。 It is necessary to perform hot rolling with a cumulative reduction rate of 40% or more in the non-recrystallization temperature range. The increase in the amount of reduction in the non-recrystallization temperature region contributes to the refinement of austenite grains during rolling, and as a result, has the effect of refining the ferrite grains and improving the mechanical properties. In order to obtain such an effect, the cumulative rolling reduction in the non-recrystallized region needs to be 40% or more. For this reason, the cumulative reduction amount in the non-recrystallized region is limited to 40% or more.
 鋼片は800~700℃の温度領域で熱間圧延を完了させた後、空冷する必要がある。この場合、冷却速度5℃/s未満で緩冷却するのが望ましい。本発明では800~700℃の二相域温度で圧延を完了し、フェライトとパーライトの混合した組織を生成する。これによって、DWTTなどの母材靭性と高強度、高延性がはかれる。なお、化学成分にもよるが強度、延性、靭性のバランス確保の観点から、熱間圧延が780~720℃の温度領域であるとさらに望ましい。 The steel slab needs to be air-cooled after hot rolling is completed in the temperature range of 800 to 700 ° C. In this case, it is desirable to cool slowly at a cooling rate of less than 5 ° C./s. In the present invention, rolling is completed at a two-phase region temperature of 800 to 700 ° C. to produce a mixed structure of ferrite and pearlite. Thereby, base material toughness such as DWTT, high strength, and high ductility are achieved. Although depending on chemical components, from the viewpoint of securing a balance of strength, ductility, and toughness, it is more desirable that the hot rolling is in a temperature range of 780 to 720 ° C.
 圧延終了温度が800℃を超えるとバンド状のパーライト組織が形成されず延性や母材靭性が低下する。また、700℃未満となると加工フェライト量が増加し延性(一様伸び)を著しく低下させる。
 本発明では冷却方法を空冷とのみ規定し、冷却速度は板厚との関係において定める。一般的な板厚では、冷却速度5℃/sを越すとMAやベイナイトが生成しやすくなり、靭性や延性を低下させる。このため、冷却速度5℃/s未満での緩冷却が望ましい。更に望ましくは2℃/s以下がよい。空冷を行うことによって容易に上記のような冷却速度を得る事ができる。
When the rolling end temperature exceeds 800 ° C., a band-like pearlite structure is not formed, and ductility and base material toughness are lowered. Moreover, when it becomes less than 700 degreeC, the amount of work ferrite will increase and ductility (uniform elongation) will fall remarkably.
In the present invention, the cooling method is defined only as air cooling, and the cooling rate is determined in relation to the plate thickness. With a general plate thickness, when the cooling rate exceeds 5 ° C./s, MA and bainite are likely to be formed, and the toughness and ductility are reduced. For this reason, slow cooling at a cooling rate of less than 5 ° C./s is desirable. More preferably, it is 2 ° C./s or less. By performing air cooling, the above cooling rate can be easily obtained.
 冷却後に、500~300℃の温度領域で温焼戻し処理を施すことは有効である。500℃超の焼戻温度では強度の低下をまねき、300℃未満ではMAが分解しないため延性の低下をまねく。また、500~300℃の温度領域で温焼戻し処理は脱水素の観点からも局部伸びを向上させる。 It is effective to perform a tempering treatment in the temperature range of 500 to 300 ° C. after cooling. If the tempering temperature exceeds 500 ° C., the strength is lowered, and if it is less than 300 ° C., the MA is not decomposed, so that the ductility is lowered. Further, the thermal tempering treatment in the temperature range of 500 to 300 ° C. improves local elongation from the viewpoint of dehydrogenation.
 なお、本発明では組織としては前述のようにフェライトとパーライトの二相組織とする必要がある。組織分率については必ずしも規定しないが、望ましくは、フェライト分率は60~95%程度が良い。 In the present invention, the structure needs to be a two-phase structure of ferrite and pearlite as described above. The structure fraction is not necessarily specified, but preferably the ferrite fraction is about 60 to 95%.
 上記でも記載したように硬化した組織が多量に入ると局部伸びを低下させる。このため本発明では、ベイナイトの生成を回避する。本発明の条件で製造した鋼板は、光学顕微鏡で測定したときに、ベイナイトが検出されない。この結果、本願の発明鋼はフェライトとパーライトの二相組織を有する。しかしながら、工業的には完全にベイナイトを除去することは困難であり、本発明でも電子顕微鏡レベルではベイナイト組織が確認されている。組織分率としては、望ましくは、ベイナイトを1%以下にするとよい。 As described above, when a hardened structure enters a large amount, local elongation is reduced. For this reason, in this invention, the production | generation of bainite is avoided. In the steel sheet produced under the conditions of the present invention, bainite is not detected when measured with an optical microscope. As a result, the present invention steel has a two-phase structure of ferrite and pearlite. However, industrially, it is difficult to completely remove bainite, and in the present invention, a bainite structure is confirmed at the electron microscope level. As the tissue fraction, bainite is desirably 1% or less.
 本発明では板厚を18mm以上、降伏強度を450MPa以上とした。板厚18mm未満では降伏強度の確保は容易であるが、空冷でも冷却速度が速くなるためベイナイトなどの硬化組織が多量に生成し伸びの確保が困難である。また、一般的には降伏強度の低下に伴い伸びが向上することが知られている。540MPa未満では化学成分、組織や製造方法などを制御しなくても、全伸びをGOST引張試験で20%以上、API引張試験で40%以上の成績を容易に得られる。このため本発明は、降伏強度が450MPa以上の鋼板において特に好適な作用効果を発揮する。さらに好ましくは、降伏強度が540MPa以上の鋼板であってもよい。鋼板の靱性はDWTT試験による-20℃の延性破面率で70%以上が好ましい。 In the present invention, the plate thickness is 18 mm or more, and the yield strength is 450 MPa or more. When the plate thickness is less than 18 mm, it is easy to ensure the yield strength. However, even with air cooling, the cooling rate increases, so that a large amount of a hardened structure such as bainite is generated and it is difficult to ensure the elongation. In general, it is known that the elongation improves as the yield strength decreases. If it is less than 540 MPa, it is possible to easily obtain a total elongation of 20% or more in the GOST tensile test and 40% or more in the API tensile test without controlling the chemical composition, structure, manufacturing method and the like. For this reason, this invention exhibits an especially suitable effect in the steel plate whose yield strength is 450 Mpa or more. More preferably, the steel sheet may have a yield strength of 540 MPa or more. The toughness of the steel sheet is preferably 70% or more at a ductile fracture surface ratio of −20 ° C. by the DWTT test.
 次に、本発明の実施例について述べる。 Next, examples of the present invention will be described.
 表1の化学成分を有する溶鋼を鋳造したスラブを、表2に示す条件にて熱間圧延を行い鋼板とした。なお、表2において、「累積圧下量」の欄は850℃以上の温度域における累積圧下率を示す。
 この後、機械的性質を評価するために試験を実施した。表2の冷却方法欄において、比較鋼r、sでは括弧に示す速度(℃/s)で加速冷却した。上記以外の鋼板は空冷で冷却された。
A slab obtained by casting molten steel having the chemical components shown in Table 1 was hot-rolled under the conditions shown in Table 2 to obtain a steel plate. In Table 2, the “cumulative reduction amount” column indicates the cumulative reduction rate in a temperature range of 850 ° C. or higher.
This was followed by a test to evaluate the mechanical properties. In the cooling method column of Table 2, the comparative steels r and s were accelerated and cooled at the rate indicated in parentheses (° C./s). The steel plates other than the above were cooled by air cooling.
 引張試験片のために、各サンプル鋼板からロシア規格のGOST試験片を採取し、YS(降伏強度:0.5%アンダーロード)、TS(引張強度)および全伸び(Total Elongation:T.El)、一様伸び(Uniform Elongation:U.El)、局部伸び(Local Elongation:L.El)を評価した。母材靱性についてはDWTT試験にて-20℃の延性破面率を評価し、70%以上を合格と判定した。 For tensile test specimens, Russian standard GOST test specimens were collected from each sample steel plate, and YS (yield strength: 0.5% underload), TS (tensile strength), and total elongation (Total Elongation: T.El). The uniform elongation (Uniform Elongation: U. El) and the local elongation (Local Elongation: L. El) were evaluated. As for the base metal toughness, the ductile fracture surface ratio at −20 ° C. was evaluated by the DWTT test, and 70% or more was determined to be acceptable.
 表3は、各鋼における機械的性質をまとめたものを示す。鋼a~oは本発明の実施例である。表1および表2から明らかなようにこれらの鋼板は化学成分と製造条件の各要件を満足しており、表3に示すように、母材強度、延性や靭性は良好である。なお、鋼組織は全てフェライト+パーライトの二相組織であり(光学顕微鏡下)、MAも1.5%以下であった。 Table 3 shows a summary of the mechanical properties of each steel. Steels a to o are examples of the present invention. As is apparent from Tables 1 and 2, these steel sheets satisfy the requirements of chemical components and production conditions, and as shown in Table 3, the base metal strength, ductility and toughness are good. The steel structure was a ferrite + pearlite two-phase structure (under an optical microscope), and MA was 1.5% or less.
 これに対し、鋼p~adは本願発明の範囲を逸脱する比較鋼である。鋼p~v、ab~adまでは製造条件が本願発明と異なり、鋼w~aaは化学成分が本発明の範囲外である。この結果、鋼p~adでは、母材の機械的性質の一つまたは複数の点で本願発明鋼に劣っていた。 In contrast, steels p to ad are comparative steels that depart from the scope of the present invention. The production conditions for steels p to v and ab to ad are different from those of the present invention, and the chemical components of steels w to aa are outside the scope of the present invention. As a result, the steels p to ad were inferior to the steel of the present invention in one or more of the mechanical properties of the base material.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
  表3中の「組織」欄は鋼板中の組織を示し、Fがフェライト、Pがパーライト、Bがベイナイトを表す。
Figure JPOXMLDOC01-appb-T000003
The “structure” column in Table 3 indicates the structure in the steel sheet, where F is ferrite, P is pearlite, and B is bainite.
 表1~3を参照してわかるように、鋼pは累積圧下量が低いため靭性が低下した。鋼qは圧延終了温度が高く、組織制御ができないため靭性が低下した。鋼r、鋼sは加速冷却のため、その過程でベイナイトやMAが多量に生成したために延性が低下した。鋼rは焼戻しによりMAが分解しているがベイナイトがあるため、靭性は回復するが伸びの回復は不十分である。鋼t、ab~adは、スラブ製造段階の軽圧下処理を省略することによって板の偏析度を上げて故意にMAを多く生成させた。そのために伸びが低下した例である。鋼uでは板厚が薄く冷却速度が高いため焼戻し前ではそもそもMAの生成が多い。鋼uはその後焼き戻しをするが温度が低いためMAが分解されず伸びと靭性が得られない。鋼vは焼戻し処理温度が高く、MA量は低減されたが、降伏強度が低い。
 鋼wはC量が低いため母材強度が低下した。また、鋼xはSi量が高いためMAが増加し延性が低下した。鋼yはMn量が高いため、MAの増加とMAそのものが硬化するため所定の伸び特性、靭性が得られない。鋼zはAl量が少ないため弱脱酸である。また鋼aaはCa量が高い。このため、鋼z、aaでは比較的粗大な酸化物が生成し、十分な伸びが得られない。
As can be seen with reference to Tables 1 to 3, steel p had low toughness due to low cumulative reduction. Steel q had a high rolling end temperature and was unable to control the structure, so the toughness decreased. Since the steel r and steel s were accelerated and cooled, a large amount of bainite and MA were produced in the process, resulting in a decrease in ductility. In steel r, MA is decomposed by tempering, but because there is bainite, the toughness is recovered, but the recovery of elongation is insufficient. Steels t, ab to ad intentionally produced a large amount of MA by increasing the segregation degree of the plate by omitting the light reduction treatment in the slab manufacturing stage. For this reason, the elongation is reduced. Since steel u has a small plate thickness and a high cooling rate, MA is primarily generated before tempering. Steel u is subsequently tempered, but since the temperature is low, MA is not decomposed and elongation and toughness cannot be obtained. Steel v has a high tempering treatment temperature, and the amount of MA is reduced, but the yield strength is low.
Since the steel C has a low C content, the strength of the base material has decreased. Moreover, since steel x has a high Si content, MA increased and ductility decreased. Since the steel y has a high Mn content, the increase in MA and the MA itself harden, so that predetermined elongation characteristics and toughness cannot be obtained. Steel z is weakly deoxidized because of its low Al content. Steel aa has a high Ca content. For this reason, a relatively coarse oxide is produced in steels z and aa, and sufficient elongation cannot be obtained.
 また、比較鋼p~vでは、本発明の組成を満足する鋳片1~7を使用しているが、製造条件が本発明と異なるため、伸びおよび/または靭性が劣っている。 Further, in the comparative steels p to v, slabs 1 to 7 satisfying the composition of the present invention are used, but the elongation and / or toughness is inferior because the manufacturing conditions are different from those of the present invention.
 以上述べたように、本発明によれば強度、延性の良好なラインパイプ用鋼が得られるため、産業上極めて有用である。 As described above, according to the present invention, steel for line pipes having good strength and ductility can be obtained, which is extremely useful industrially.

Claims (11)

  1.  質量%で、
      C:0.07~0.15%、
      Si:0.05~0.60%、
      Mn:0.80~1.80%、
      P:0.010以下、
      S:0.007%以下、
      V:0.05~0.12%、
      Nb: 0.005~0.070%、
      Al:0.005~0.08%、
      Ti:0.005~0.030%、
      Ca:0.0005~0.0035%、
      N:0.0020~0.0060%、
      O:0.0030%以下を含有し、残部が鉄および不可避的不純物からなり、
     組織がフェライトとパーライトの二相組織であり、
     島状マルテンサイトの面積分率が1.5%未満であり、
     板厚が18mm以上であり、降伏強度が450MPa以上である
    ことを特徴とする、強度および延性の良好なラインパイプ用鋼。
    % By mass
    C: 0.07 to 0.15%,
    Si: 0.05 to 0.60%,
    Mn: 0.80 to 1.80%,
    P: 0.010 or less,
    S: 0.007% or less,
    V: 0.05 to 0.12%,
    Nb: 0.005 to 0.070%,
    Al: 0.005 to 0.08%,
    Ti: 0.005 to 0.030%,
    Ca: 0.0005 to 0.0035%,
    N: 0.0020 to 0.0060%,
    O: contains 0.0030% or less, the balance consists of iron and inevitable impurities,
    The structure is a two-phase structure of ferrite and pearlite,
    The area fraction of island martensite is less than 1.5%,
    A steel for line pipes with good strength and ductility, characterized in that the plate thickness is 18 mm or more and the yield strength is 450 MPa or more.
  2.  更に、質量%で、
      Cu:0.05~0.70%、
      Ni:0.05~0.70%、
      Cr:0.80%以下、
      Mo:0.30%以下、
      B:0.0003~0.0030%、
      Mg:0.0003~0.0030%、
      REM:0.0005~0.0050%の一種または二種以上を含有したことを特徴とする請求項1に記載の強度および延性の良好なラインパイプ用鋼。
    Furthermore, in mass%,
    Cu: 0.05 to 0.70%,
    Ni: 0.05 to 0.70%,
    Cr: 0.80% or less,
    Mo: 0.30% or less,
    B: 0.0003 to 0.0030%,
    Mg: 0.0003 to 0.0030%,
    The steel for line pipes with good strength and ductility according to claim 1, characterized by containing one or more of REM: 0.0005 to 0.0050%.
  3.  光学顕微鏡下でベイナイト組織が検出されないことを特徴とする請求項1または2に記載の強度および延性の良好なラインパイプ用鋼。 3. Steel for line pipes with good strength and ductility according to claim 1 or 2, wherein a bainite structure is not detected under an optical microscope.
  4.  Mnの偏析度が1.7以下であることを特徴とする請求項1または2に記載の強度および延性の良好なラインパイプ用鋼。 3. The steel for line pipes with good strength and ductility according to claim 1 or 2, wherein the segregation degree of Mn is 1.7 or less.
  5.  Si,Pの偏析度がそれぞれ1.5以下、8.0以下であることを特徴とする請求項4に記載の強度および延性の良好なラインパイプ用鋼。 The segregation degree of Si and P is 1.5 or less and 8.0 or less, respectively, and the steel for line pipes having good strength and ductility according to claim 4.
  6.  請求項1または2に記載の化学成分の鋳片を1250℃以下の温度に加熱後、850℃以上の温度域において累積圧下率で40%以上の熱間圧延をし、700~800℃の温度領域で熱間圧延を終了させた後、空冷することを特徴とする強度および延性の良好なラインパイプ用鋼の製造方法。 3. The slab of the chemical component according to claim 1 or 2 is heated to a temperature of 1250 ° C. or lower, and then hot-rolled at a cumulative reduction ratio of 40% or higher in a temperature range of 850 ° C. or higher to a temperature of 700 to 800 ° C. A method for producing steel for line pipes with good strength and ductility, characterized by air cooling after finishing hot rolling in a region.
  7.  前記空冷後、鋼板に500~300℃の温度領域で焼戻し処理を施すことを特徴とする請求項6に記載の強度および延性の良好なラインパイプ用鋼の製造方法。 The method for producing steel for a line pipe with good strength and ductility according to claim 6, wherein after the air cooling, the steel sheet is tempered in a temperature range of 500 to 300 ° C.
  8.  熱間圧延によって製造された鋼板中の島状マルテンサイト量を測定し、
     島状マルテンサイトの面積分率が1.5%未満になるように製造条件を制御することによって延性を改善させることを特徴とする強度および延性の良好なラインパイプ用鋼の品質管理方法。
    Measure the amount of island martensite in the steel sheet manufactured by hot rolling,
    A quality control method for steel for line pipes having good strength and ductility, wherein ductility is improved by controlling production conditions so that the area fraction of island martensite is less than 1.5%.
  9.  熱間圧延後の冷却方法を空冷とすることによって延性を改善させることを特徴とする請求項8に記載の強度および延性の良好なラインパイプ用鋼の品質管理方法。 The quality control method for steel for line pipes with good strength and ductility according to claim 8, wherein the ductility is improved by air cooling as a cooling method after hot rolling.
  10.  鋼板段階のMnの偏析度を1.7以下とすることによって延性を改善させることを特徴とする請求項8または9に記載の強度および延性の良好なラインパイプ用鋼の品質管理方法。 The method for quality control of steel for line pipes with good strength and ductility according to claim 8 or 9, wherein the ductility is improved by setting the segregation degree of Mn at a steel plate stage to 1.7 or less.
  11.  鋼板段階のベイナイト組織分率を1%以下にすることによって延性を改善させることを特徴とする請求項8または9に記載の強度および延性の良好なラインパイプ用鋼の品質管理方法。 The method for quality control of steel for line pipes with good strength and ductility according to claim 8 or 9, wherein the ductility is improved by setting the bainite structure fraction in the steel plate stage to 1% or less.
PCT/JP2010/067351 2009-10-05 2010-10-04 Steel for linepipe having good strength and malleability, and method for producing the same WO2011043287A1 (en)

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JP2013019014A (en) * 2011-07-11 2013-01-31 Jfe Steel Corp Steel for weld structure superior in ctod property at large heat input weld heat affected zone, and production method thereof
CN113025885A (en) * 2021-02-08 2021-06-25 江阴兴澄特种钢铁有限公司 Low-yield-ratio high-strength pipeline steel plate with good HIC (hydrogen induced cracking) resistance and manufacturing method thereof
CN114737120A (en) * 2022-04-02 2022-07-12 鞍钢股份有限公司 Steel for large-diameter tube bundle outer bearing tube and preparation method thereof

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JP2008101242A (en) * 2006-10-19 2008-05-01 Jfe Steel Kk High-strength steel plate with excellent hic resistance for line pipe, and its manufacturing method
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JP2008101242A (en) * 2006-10-19 2008-05-01 Jfe Steel Kk High-strength steel plate with excellent hic resistance for line pipe, and its manufacturing method
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Publication number Priority date Publication date Assignee Title
JP2013019014A (en) * 2011-07-11 2013-01-31 Jfe Steel Corp Steel for weld structure superior in ctod property at large heat input weld heat affected zone, and production method thereof
CN113025885A (en) * 2021-02-08 2021-06-25 江阴兴澄特种钢铁有限公司 Low-yield-ratio high-strength pipeline steel plate with good HIC (hydrogen induced cracking) resistance and manufacturing method thereof
CN114737120A (en) * 2022-04-02 2022-07-12 鞍钢股份有限公司 Steel for large-diameter tube bundle outer bearing tube and preparation method thereof
CN114737120B (en) * 2022-04-02 2022-11-18 鞍钢股份有限公司 Steel for large-diameter tube bundle outer bearing tube and preparation method thereof

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