JP4940886B2 - High strength steel plate for line pipe with excellent HIC resistance and method for producing the same - Google Patents
High strength steel plate for line pipe with excellent HIC resistance and method for producing the same Download PDFInfo
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- 229910000831 Steel Inorganic materials 0.000 title claims description 86
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- 229910052750 molybdenum Inorganic materials 0.000 claims description 35
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- 229910052758 niobium Inorganic materials 0.000 claims description 21
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- 229910052748 manganese Inorganic materials 0.000 claims description 5
- 229910052791 calcium Inorganic materials 0.000 claims description 4
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- RWSOTUBLDIXVET-UHFFFAOYSA-N Dihydrogen sulfide Chemical compound S RWSOTUBLDIXVET-UHFFFAOYSA-N 0.000 description 1
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Description
本発明は、ＡＰＩ規格Ｘ７０グレード以上の強度を有する高強度鋼板およびその製造方法に関し、特に耐水素誘起割れ性（耐ＨＩＣ性）に優れたものに関する。 The present invention relates to a high-strength steel plate having a strength of API standard X70 grade or higher and a method for producing the same, and particularly to a material excellent in hydrogen-induced crack resistance (HIC resistance).
硫化水素を含む原油や天然ガスの輸送に用いられるラインパイプは、強度、靭性、溶接性の他に、耐水素誘起割れ性（以下、耐ＨＩＣ性）や耐応力腐食割れ性（以下、耐ＳＣＣ性）などのいわゆる耐サワー性が必要とされる。 In addition to strength, toughness and weldability, line pipes used for transporting crude oil and natural gas containing hydrogen sulfide are resistant to hydrogen-induced cracking resistance (hereinafter referred to as HIC resistance) and stress corrosion cracking resistance (hereinafter referred to as SCC resistance). So-called sour resistance is required.
鋼材の水素誘起割れ（以下、ＨＩＣ）は、腐食反応による水素イオンが鋼材表面に吸着し、原子状の水素として鋼内部に侵入、鋼中のＭｎＳなどの非金属介在物や硬い第２相組織のまわりに拡散・集積し、その内圧により割れを生ずるとされている。 Hydrogen-induced cracking (hereinafter referred to as HIC) in steel materials is that hydrogen ions from the corrosion reaction are adsorbed on the steel material surface and penetrate into the steel as atomic hydrogen, and non-metallic inclusions such as MnS in the steel and a hard second phase structure. It is said that it diffuses and accumulates around the surface and cracks occur due to its internal pressure.
特許文献１には、ＣａやＣｅをＳ量に対して適量添加することにより、針状のＭｎＳの生成を抑制し、応力集中の小さい微細に分散した球状の介在物に形態を変えて割れの発生・伝播を抑制する、耐ＨＩＣ性の優れたラインパイプ用鋼の製造方法が開示されている。 In Patent Document 1, by adding an appropriate amount of Ca or Ce to the amount of S, the formation of acicular MnS is suppressed, and the shape is changed to a finely dispersed spherical inclusion with a small stress concentration, and cracks are not generated. A method for producing steel for line pipes that suppresses generation and propagation and has excellent HIC resistance is disclosed.
特許文献２，３には、偏析傾向の高い元素（Ｃ、Ｍｎ、Ｐ等）の低減や、スラブ加熱段階での均熱処理、冷却時の変態途中での加速冷却により、中心偏析部での割れの起点となる島状マルテンサイト、割れの伝播経路となるマルテンサイトやベイナイトなどの硬化組織の生成を抑制した、耐ＨＩＣ性に優れた鋼が開示されている。
In
特許文献１〜３に記載の耐ＨＩＣ性を改善する方法はいずれも中心偏析部が対象であり、加速冷却または直接焼入れによって製造され、冷却速度の速い鋼板表面部が内部に比べ硬化し、表面近傍からも水素誘起割れが発生するＡＰＩ Ｘ７０グレードを超える高強度鋼板を対象としたものではない。 The methods for improving the HIC resistance described in Patent Documents 1 to 3 are all about the center segregation part, manufactured by accelerated cooling or direct quenching, and the steel plate surface part having a high cooling rate is hardened compared to the inside, and the surface It is not intended for high-strength steel sheets exceeding the API X70 grade that generate hydrogen-induced cracks from the vicinity.
加速冷却によって得られるこれらの高強度鋼板のミクロ組織は、表面のみならず内部までベイナイトまたはアシキュラーフェライトの比較的割れ感受性の高い組織となるため、耐ＨＩＣ性に対しては、更なる厳格な偏析抑制が必要となるだけでなく、硫化物系または酸化物系介在物を起点としたＨＩＣへの対策が必要である。 Since the microstructure of these high-strength steel sheets obtained by accelerated cooling is a relatively high cracking susceptibility of bainite or acicular ferrite not only to the surface but also to the inside, it is even more stringent for HIC resistance. Not only segregation suppression is required, but also countermeasures against HIC starting from sulfide-based or oxide-based inclusions are necessary.
すなわち、ＡＰＩ Ｘ７０グレードを超える高強度鋼板の場合、強度を確保しつつも更なる厳格な偏析抑制を可能とする成分設計・ミクロ組織制御に加えて、鋼板の表面部のＨＩＣまたは、硫化物系や酸化物系介在物を起点としたＨＩＣへの対策が必要である。 That is, in the case of high-strength steel sheets exceeding API X70 grade, in addition to the component design and microstructure control that enables further strict segregation control while ensuring the strength, the HIC or sulfide system on the surface of the steel sheet And measures against HIC starting from oxide inclusions are necessary.
特許文献４には、ミクロ組織が、割れ感受性の高いブロック状ベイナイトやマルテンサイトを含まない耐ＨＩＣ性に優れた高強度鋼として、フェライト−ベイナイト２相組織である、ＡＰＩ Ｘ８０グレードの耐ＨＩＣ性に優れた高強度鋼材が開示されている。
また、特許文献５，６には、ミクロ組織をフェライト単相組織とすることで耐ＳＣＣ（ＳＳＣＣ）性や耐ＨＩＣ性を改善し、強度は、ＭｏまたはＴｉの多量添加によって得られる炭化物の析出強化で確保する高強度鋼が開示されている。
しかし、特許文献４に記載の高強度鋼のベイナイト組織は、ブロック状ベイナイトやマルテンサイト程ではないが比較的割れ感受性の高い組織であり、ＳおよびＭｎ量を厳しく制限して、Ｃａ処理を必須として耐ＨＩＣ性を向上させる必要があるため、製造コストが高くなることが懸念される。また、特許文献４に記載の圧延・冷却方法を用いてフェライト−ベイナイト２相組織を安定的に得ることは難しい。
However, the bainite structure of the high-strength steel described in
一方、特許文献５，６に記載のフェライト相は延性に富んだ組織であり、割れ感受性が極めて低いため、ベイナイト組織またはアシキュラーフェライト組織の鋼に比べ耐ＨＩＣ性が大幅に改善される。
On the other hand, the ferrite phases described in
しかし、フェライト単相では強度が低いため、特許文献５に記載の鋼はＣ及びＭｏを多量に添加した鋼を用いて、炭化物を多量に析出させることによって高強度化し、特許文献６記載の鋼帯ではＴｉ添加鋼を特定の温度で鋼帯に巻き取り、ＴｉＣの析出強化を利用して高強度化を図っている。
However, since the strength of the ferrite single phase is low, the steel described in
ところが、特許文献５に記載の、Ｍｏ炭化物が分散したフェライト組織を得るためには、焼入れ焼戻しの後に冷間加工を行い、さらに再度焼戻しを行う必要があり、製造コストが上昇する。
However, in order to obtain a ferrite structure in which Mo carbide is dispersed as described in
また、Ｍｏ炭化物の粒径が約０．１ミクロンと大きく、強度上昇効果が低いため、Ｃ及びＭｏの含有量を高め、炭化物の量をふやすことによって所定の強度を得る必要があり、高価な成分組成となる。 In addition, since the particle size of Mo carbide is as large as about 0.1 microns and the effect of increasing the strength is low, it is necessary to increase the content of C and Mo and obtain a predetermined strength by increasing the amount of carbide, which is expensive. It becomes an ingredient composition.
特許文献６に記載の高強度鋼で、強度確保のために利用するＴｉＣは、Ｍｏ炭化物に比べ微細で、析出強化に有効であるが、析出時の温度の影響を受けて粗大化しやすい。特許文献６では粗大化対策が十分でないため析出強化が小さく、多量のＴｉ添加が必要となっている。
TiC used for securing strength in the high-strength steel described in
そこで、本発明は、上述した課題を解決し、ＡＰＩ Ｘ７０グレード以上のラインパイプ用高強度鋼板であって、中央偏析部及び表面近傍や介在物周辺において、優れた耐ＨＩＣ特性を有するラインパイプ用高強度鋼板およびその製造方法を提供することを目的とする。 Therefore, the present invention solves the above-described problems, and is a high-strength steel sheet for API X70 grade linepipe, which has excellent HIC resistance near the central segregation part and the surface and inclusions. An object is to provide a high-strength steel sheet and a method for producing the same.
本発明者らは耐ＨＩＣ特性向上と高強度の両立のために、鋼材の成分とミクロ組織および鋼板の製造方法を鋭意検討し、以下の知見を得た。
1．高強度と耐ＨＩＣ特性の両立には、偏析が抑制され、且つ高強度の確保が可能な成分組成とすることが必要で、偏析が比較的小さなＣｒ、ＣｕおよびＮｉを含有し、Ｃｅｑ（以下、Ｃｅｑ１）および偏析を考慮したＣｅｑ（以下、Ｃｅｑ２）を一定の範囲とし、析出強化が得られる系とする。
２．また、ミクロ組織を、フェライト組織とベイナイト組織との強度差の小さい、フェライト＋ベイナイト２相組織とすることが、最も効果的である。
The present inventors diligently studied the components and microstructure of the steel material and the manufacturing method of the steel sheet in order to achieve both high HIC resistance and high strength, and obtained the following knowledge.
1． In order to achieve both high strength and HIC resistance, it is necessary to have a component composition that can suppress segregation and ensure high strength, and contains Cr, Cu, and Ni that have relatively small segregation. , Ceq1) and Ceq in consideration of segregation (hereinafter, Ceq2) are set within a certain range, and a system in which precipitation strengthening is obtained is obtained.
2. It is most effective to make the microstructure a ferrite + bainite two-phase structure with a small strength difference between the ferrite structure and the bainite structure.
すなわち、フェライトとベイナイトの２相組織は、一般的には軟質な相（フェライト相）と硬質な相（ベイナイト相）の混合組織であり、このような組織を有する鋼材は軟質な相（フェライト相）と硬質な相（ベイナイト相）との界面に水素が集積しやすく、更に前記界面が割れの伝播経路となるため、耐ＨＩＣ特性が劣る。しかし、両者の強度差を小さくすることで、耐ＨＩＣ特性と高強度の両立が可能である。
３．上記成分組成において析出強化に寄与する成分として、Ｍｏ，Ｔｉ、更には（Ｎｂ、Ｖ）の一種または二種を複合添加し、Ｎに対するＴｉの添加量を適正化し、かつＣに対するＭｏ、Ｔｉの添加量を適正化した系とすると、炭化物による析出強化を最大限に活用することができる。
４．熱間圧延後の加速冷却とその後の再加熱で、Ｔｉ、Ｍｏ、（Ｎｂ、Ｖ）の一種または二種を含む析出物が分散析出し、フェライト相の高強度化が達成でき、硬質相であるベイナイト相の軟化が起こり、強度差の小さいフェライト＋ベイナイト２相組織を得ることができるとともに、ＭＡの生成を低減することができる。
That is, the two-phase structure of ferrite and bainite is generally a mixed structure of a soft phase (ferrite phase) and a hard phase (bainite phase), and a steel material having such a structure is a soft phase (ferrite phase). ) And a hard phase (bainite phase), hydrogen tends to accumulate, and the interface serves as a crack propagation path, resulting in poor HIC resistance. However, by reducing the difference in strength between the two, it is possible to achieve both HIC resistance and high strength.
3. As a component contributing to precipitation strengthening in the above component composition, Mo or Ti, and further, one or two of (Nb, V) are added in combination, the amount of Ti added to N is optimized, and Mo and Ti to C are added. If the addition amount is optimized, precipitation strengthening due to carbide can be utilized to the maximum.
4). With accelerated cooling after hot rolling and subsequent reheating, precipitates containing one or two of Ti, Mo, (Nb, V) are dispersed and precipitated, and the strength of the ferrite phase can be increased. Softening of a certain bainite phase occurs, and a ferrite + bainite two-phase structure with a small strength difference can be obtained, and the formation of MA can be reduced.
本発明は得られた知見を基に更に検討を加えてなされたもので、すなわち、本発明は、
１．鋼組成が、質量％で、Ｃ：０．０２〜０．０８％、Ｓｉ：０．０１〜０．５％、Ｍｎ：
０．５〜１．８％、Ｐ：０．０１％以下、Ｓ：０．００２％以下、Ｍｏ：０．０５〜０．
５％、Ｃｒ：０．０５〜１％、Ａｌ：０．０７％以下、Ｎｂ：０．００５〜０．０７％、Ｔｉ：０．０１〜０．０５％、Ｃａ：０．０００５〜０．００５％、Ｎ：０．００７％以下、下記（１）〜（４）式を満足し、残部Ｆｅおよび不可避的不純物で
金属組織がフェライトとベイナイトの２相組織であり、島状マルテンサイト（ＭＡ）の分
率が体積分率で５％以下であり、フェライト相に、ＴｉとＭｏを含む複合析出物が分散析出していることを
特徴とする、耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板。
０．３５≦Ｃｅｑ１≦０．４５・・・（１）
但し、Ｃｅｑ１＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５
で、各元素は含有量（質量％）で含有しない元素は０とする。
Ｃｅｑ２≦０．６・・・（２）
但し、Ｃｅｑ２＝１．５Ｃ＋１．６Ｍｎ／６＋（１．３Ｃｕ＋１．３Ｎｉ）／１５＋（１
．１Ｃｒ＋１．２Ｍｏ＋Ｖ）／５
但し、各元素は含有量（質量％）で含有しない元素は０とする。
Ｔｉ／Ｎ＞４・・・（３）
但し、Ｔｉ，Ｎは質量％とする。
０．５≦Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）≦３・・・（４）
但し、各元素は原子％で含有しない元素は０とする。
２．鋼組成が、更に、質量％で、Ｖ：０．００５〜０．１％を含有し、複合析出物が、Ｔｉ、Ｍｏ、および（Ｎｂ、Ｖ）の一種または二種を含むことを特徴とする、１に記載の耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板。
３．鋼組成が、更に、質量％で、Ｃｕ：０．５％以下、Ｎｉ：０．５％以下の中から選ばれる一種又は二種を含有することを特徴とする１または２に記載の耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板。
４．複合析出物が１０ｎｍ未満の径を有する微細析出物であることを特徴とする１乃至３の何れか一つに記載の耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板。
５．１〜３のいずれか一つに記載の化学成分を含有する鋼を、加熱温度：１０００〜１３００℃、圧延終了温度：Ａｒ３温度以上の条件で熱間圧延した後、冷却速度：５℃／ｓ以上で４００〜６００℃まで加速冷却を行い、冷却後直ちに昇温速度：０．５℃／ｓ以上で６００〜７００℃の温度まで再加熱を行うことを特徴とする、耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板の製造方法。
６．１〜４のいずれか一つに記載の鋼板を用いて製造されたことを特徴とする、耐ＨＩＣ特性に優れたラインパイプ。
The present invention was made by further study based on the obtained knowledge, that is, the present invention is
1. Steel composition is mass%, C: 0.02-0.08%, Si: 0.01-0.5%, Mn:
0.5-1.8%, P: 0.01% or less, S: 0.002% or less, Mo: 0.05-0.
5%, Cr: 0.05-1%, Al: 0.07% or less, Nb: 0.005-0.07%, Ti: 0.01-0.05%, Ca: 0.0005-0. 005%, N: 0.007% or less, satisfying the following formulas (1) to (4), the balance Fe and unavoidable impurities, the metal structure is a two-phase structure of ferrite and bainite, and island martensite (MA ) With a volume fraction of 5% or less, and composite precipitates containing Ti and Mo are dispersed and precipitated in the ferrite phase . High strength for line pipes with excellent HIC resistance steel sheet.
0.35 ≦ Ceq1 ≦ 0.45 (1)
However, Ceq1 = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5
Thus, each element is contained in an amount (mass%), and elements not contained are set to 0.
Ceq2 ≦ 0.6 (2)
However, Ceq2 = 1.5C + 1.6Mn / 6 + (1.3Cu + 1.3Ni) / 15 + (1
. 1Cr + 1.2Mo + V) / 5
However, each element is contained (mass%), and elements not contained are set to 0.
Ti / N> 4 (3)
However, Ti and N are mass%.
0.5 ≦ C / (Mo + Ti + Nb + V) ≦ 3 (4)
However, the element which does not contain each element in atomic% is set to 0.
2. The steel composition further includes, in mass%, V: 0.005 to 0.1%, and the composite precipitate includes one or two of Ti, Mo, and (Nb, V). The high-strength steel sheet for line pipes having excellent HIC resistance according to 1.
3. The steel composition further contains one or two kinds selected from Cu: 0.5% or less and Ni: 0.5% or less in terms of mass%, wherein the HIC resistance according to 1 or 2 High-strength steel sheet for line pipes with excellent characteristics.
4). The high strength steel sheet for a line pipe excellent in HIC resistance according to any one of 1 to 3, wherein the composite precipitate is a fine precipitate having a diameter of less than 10 nm.
The steel containing the chemical component according to any one of 5.1 to 3 is hot-rolled at a heating temperature of 1000 to 1300 ° C. and a rolling end temperature of Ar 3 or higher, and then a cooling rate of 5 HIC resistance, characterized in that accelerated cooling is performed at 400 ° C./s or higher to 400 to 600 ° C., and immediately after cooling, the heating rate is reheated to a temperature of 600 to 700 ° C. at 0.5 ° C./s or higher. Manufacturing method for high-strength steel sheets for line pipes.
A line pipe excellent in HIC resistance, characterized by being manufactured using the steel sheet according to any one of 6.1 to 4.
本発明によれば、ＡＰＩ Ｘ７０グレード以上の高強度を有し、かつ耐ＨＩＣ性の優れた電縫鋼管、スパイラル鋼管、ＵＯＥ鋼管等の鋼管が、多量の合金元素を添加することなく製造可能で、産業上極めて有用である。 According to the present invention, steel pipes such as ERW steel pipes, spiral steel pipes and UOE steel pipes having high strength of API X70 grade or higher and excellent in HIC resistance can be manufactured without adding a large amount of alloy elements. It is extremely useful in industry.
本発明ではミクロ組織、成分組成を規定する。以下、限定理由について詳細に説明する。
［ミクロ組織］
本発明では、ミクロ組織を実質的にフェライト＋ベイナイト２相組織とする。
フェライト相は延性に富んでおり割れ感受性が低いため、高い耐ＨＩＣ特性が得れ、ベイナイト相は優れた強度靭性を有している。
In the present invention, the microstructure and component composition are defined. Hereinafter, the reason for limitation will be described in detail.
[Microstructure]
In the present invention, the microstructure is substantially a ferrite + bainite two-phase structure.
Since the ferrite phase is rich in ductility and has low cracking sensitivity, high HIC resistance can be obtained, and the bainite phase has excellent strength toughness.
ＭＡは非常に硬い硬質相であることから、母相とＭＡとの界面に水素が集積しやすく、更に、前記界面が割れの伝播経路となる可能性が高いため、ＭＡ分率の上昇とともに耐ＨＩＣ特性が急激に劣化する。 Since MA is a very hard hard phase, it is easy for hydrogen to accumulate at the interface between the parent phase and MA, and the interface is likely to become a crack propagation path. HIC characteristics deteriorate rapidly.
本発明では、熱間圧延後に加速冷却し、その後の再加熱する製造プロセスにより、ＭＡの生成を抑制し耐ＨＩＣ特性の向上を可能とする。 In the present invention, the production of MA is suppressed and the HIC resistance can be improved by a manufacturing process of accelerated cooling after hot rolling and subsequent reheating.
フェライト＋ベイナイトを備えた２相組織に、マルテンサイトやパーライトなどの異なる金属組織が混在すると、異相界面での水素集積や応力集中によってＨＩＣを生じやすくなるため、フェライト相とベイナイト相以外の組織分率は少ない程良い。 If different metal structures such as martensite and pearlite are mixed in the two-phase structure with ferrite + bainite, HIC is likely to occur due to hydrogen accumulation and stress concentration at the interface between different phases. The lower the rate, the better.
しかし、フェライト相とベイナイト相以外の組織の体積分率が低い場合は影響が無視できるため、体積分率で５％以下の他の金属組織、すなわちマルテンサイト、パーライト等を１種または２種以上含有してもよい。 However, since the influence is negligible when the volume fraction of the structure other than the ferrite phase and the bainite phase is low, one or more kinds of other metal structures with a volume fraction of 5% or less, that is, martensite, pearlite, etc. are used. You may contain.
また、ベイナイト分率は特に規定しないが、母材の靭性確保の観点から１０％以上、耐ＨＩＣ特性の観点から８０％以下とすることが好ましい。より好ましくは、２０〜６０％である。 The bainite fraction is not particularly defined, but is preferably 10% or more from the viewpoint of securing the toughness of the base material and 80% or less from the viewpoint of HIC resistance. More preferably, it is 20 to 60%.
次に、フェライト相内に分散析出する析出物について説明する。本発明に係る鋼板では、フェライト相中にＭｏとＴｉを含有する析出物が分散析出することによりフェライト相が強化され、フェライト−ベイナイト間の強度差が小さくなるため、優れた耐ＨＩＣ特性を得ることができる。 Next, the precipitate that is dispersed and precipitated in the ferrite phase will be described. In the steel sheet according to the present invention, a precipitate containing Mo and Ti is dispersed and precipitated in the ferrite phase, whereby the ferrite phase is strengthened, and the strength difference between ferrite and bainite is reduced, so that excellent HIC resistance is obtained. be able to.
当該析出物は極めて微細であるので耐ＨＩＣ特性を劣化させることはない。本発明ではＭｏとＴｉを複合添加し、ＭｏとＴｉの両者を含有する複合炭化物を鋼中に微細析出させる。ＭｏＣおよび／またはＴｉＣの析出強化の場合に比べて、より大きな強度向上効果が得られることが特徴である。 Since the precipitate is extremely fine, the HIC resistance is not deteriorated. In the present invention, Mo and Ti are added in combination, and a composite carbide containing both Mo and Ti is finely precipitated in the steel. Compared to the precipitation strengthening of MoC and / or TiC, it is characterized in that a greater strength improvement effect can be obtained.
この従来にない大きな強度向上効果は、ＭｏとＴｉの両者を含有する複合炭化物の場合、熱的に安定でかつ成長速度が遅いので、粒径が１０ｎｍ未満の極めて微細であることによるものと推察される。 This unprecedented strength improvement effect is presumed to be due to the fact that the composite carbide containing both Mo and Ti is thermally stable and has a slow growth rate, so that the particle size is extremely fine with a particle size of less than 10 nm. Is done.
尚、ＭｏとＴｉを含有する複合炭化物は、Ｍｏ、Ｔｉ、Ｃのみで構成される場合、ＭｏとＴｉの合計量と、Ｃ量とが原子比で１：１の付近で化合している。 In addition, when the composite carbide containing Mo and Ti is composed only of Mo, Ti, and C, the total amount of Mo and Ti and the amount of C are combined in the vicinity of 1: 1 by atomic ratio.
また、本発明では、溶接熱影響部の靭性をより向上させる場合、Ｎｂ、Ｖの一種または二種を複合添加するが、析出物がＭｏ、Ｔｉ、および（Ｎｂ、Ｖ）の一種または二種を含んだ複合炭化物となり、ＭｏとＴｉを含有する複合炭化物と同様の析出強化が得られる。 In the present invention, in order to further improve the toughness of the heat affected zone, one or two of Nb and V are added in combination, but the precipitate is one or two of Mo, Ti, and (Nb, V). Thus, precipitation strengthening similar to that of composite carbide containing Mo and Ti is obtained.
ＭｏとＴｉを含有する複合炭化物や更に（Ｎｂ、Ｖ）の一種または二種を含んだ複合炭化物は微細炭化物であり、主にフェライト相中に析出するが、化学成分、製造条件によってはベイナイト相からも析出する場合もある。 Composite carbides containing Mo and Ti and further composite carbides containing one or two of (Nb, V) are fine carbides, and mainly precipitate in the ferrite phase, but depending on the chemical composition and production conditions, the bainite phase It may also be deposited from
この場合は更なる強化が可能であるが、ベイナイト相が析出強化されても、フェライト相とベイナイト相の硬度差がＨＶ７０以下なら耐ＨＩＣ性能を損なうことはない。 In this case, further strengthening is possible, but even if the bainite phase is strengthened by precipitation, if the hardness difference between the ferrite phase and the bainite phase is HV70 or less, the HIC resistance is not impaired.
これら１０ｎｍ未満の析出物の個数は、降伏強度が５００ＭＰａ以上（ＡＰＩ×７０グレード以上）の高強度鋼板とするためには、２×１０３個／μｍ３以上析出させることが好ましい。析出形態は、ランダムでも列状でも良く、特に規定しない。 In order to obtain a high-strength steel sheet having a yield strength of 500 MPa or more (API × 70 grade or more), it is preferable to deposit 2 × 10 3 pieces / μm 3 or more of these precipitates of less than 10 nm. The form of precipitation may be random or in line, and is not particularly defined.
また、本発明では、ＭｏとＴｉの複合炭化物による高強度化の効果を損なわず耐ＨＩＣ特性を劣化させない程度であれば、ＭｏとＴｉを主体とする複合炭化物以外の析出物が生成しても良い。但し、１０ｎｍ未満の析出物の個数はＴｉＮを除いた全析出物の個数の９５％以上であることが好ましい。 In the present invention, even if precipitates other than the composite carbide mainly composed of Mo and Ti are generated as long as the effect of increasing the strength by the composite carbide of Mo and Ti is not impaired and the HIC resistance is not deteriorated. good. However, the number of precipitates of less than 10 nm is preferably 95% or more of the total number of precipitates excluding TiN.
［化学成分］
次に、化学成分について説明する。以下の説明において特に記載がない場合は、％で示す単位は全て質量％である。
[Chemical composition]
Next, chemical components will be described. Unless otherwise specified in the following description, all units shown in% are% by mass.
Ｃ
Ｃは０．０２〜０．０８％とする。Ｃは炭化物として析出強化に寄与する元素であるが、０．０２％未満では十分な強度が確保できず、０．０８％を超えると靭性や耐ＨＩＣ性を劣化させるため、含有量を０．０２〜０．０８％に規定する。
C
C is 0.02 to 0.08%. C is an element that contributes to precipitation strengthening as a carbide. However, if it is less than 0.02%, sufficient strength cannot be secured, and if it exceeds 0.08%, toughness and HIC resistance are deteriorated. It is specified at 02 to 0.08%.
Ｓｉ
Ｓｉは、０．０１〜０．５％とする。Ｓｉは脱酸のため添加するが、０．０１％未満では脱酸効果が十分でなく、０．５％を超えると靭性や溶接性を劣化させるため、含有量を０．０１〜０．５％に規定する。
Si
Si is set to 0.01 to 0.5%. Si is added for deoxidation, but if it is less than 0.01%, the deoxidation effect is not sufficient, and if it exceeds 0.5%, the toughness and weldability are deteriorated. %.
Ｍｎ
Ｍｎは、０．５〜１．８％とする。Ｍｎは強度、靭性のため添加するが、０．５％未満ではその効果が十分でなく、１．８％を超えると溶接性と耐ＨＩＣ性が劣化するため、Ｍｎ含有量を０．５〜１．８％に規定する。好ましくは、０．５〜１．５％である。
Mn
Mn is 0.5 to 1.8%. Mn is added for strength and toughness, but if less than 0.5%, the effect is not sufficient, and if it exceeds 1.8%, the weldability and HIC resistance deteriorate, so the Mn content is 0.5 to It is specified to 1.8%. Preferably, it is 0.5 to 1.5%.
Ｐ
Ｐは、０．０１％以下とする。Ｐは溶接性と耐ＨＩＣ性を劣化させる不可避不純物元素であるため、含有量の上限を０．０１％に規定する。
P
P is 0.01% or less. Since P is an inevitable impurity element that deteriorates weldability and HIC resistance, the upper limit of the content is specified as 0.01%.
Ｓ
Ｓは、０．００２％以下とする。Ｓは一般的には鋼中においてはＭｎＳ介在物となり耐ＨＩＣ特性を劣化させるため少ないほどよい。しかし、０．００２％以下であれば問題ないため、Ｓ含有量の上限を０．００２％に規定する。
S
S is made 0.002% or less. In general, S is preferably as small as possible because it becomes MnS inclusions in steel and deteriorates the HIC resistance. However, since there is no problem if it is 0.002% or less, the upper limit of the S content is defined as 0.002%.
Ｍｏ
Ｍｏは、０．０５〜０．５％とする。Ｍｏは本発明において重要な元素であり、０．０５％以上含有させることで、熱間圧延後冷却時のパーライト変態を抑制しつつ、Ｔｉとの微細な複合析出物を形成し、強度上昇に大きく寄与する。
Mo
Mo is set to 0.05 to 0.5%. Mo is an important element in the present invention, and by containing 0.05% or more, fine composite precipitates with Ti are formed while suppressing pearlite transformation during cooling after hot rolling, thereby increasing strength. A big contribution.
しかし、０．５％を超えて添加するとマルテンサイトなどの硬化相を形成し耐ＨＩＣ特性が劣化するため、Ｍｏ含有量を０．０５〜０．５％に規定する。好ましくは、０．０５〜０．３％未満である。 However, if added over 0.5%, a hardened phase such as martensite is formed and the HIC resistance is deteriorated, so the Mo content is specified to be 0.05 to 0.5%. Preferably, it is 0.05 to less than 0.3%.
Ｃｒ
Ｃｒは、０．０５〜１％とする。ＣｒはＭｎと同様に低Ｃでも十分な強度を得るために有効な元素であり、偏析も比較的小さいため、強度確保と耐ＨＩＣ特性向上に有効な元素である。０．０５％未満ではその効果が十分でなく、１％を超えて添加すると溶接性を劣化するため、Ｃｒ含有量を０．０５〜１％に規定する。
Cr
Cr is 0.05 to 1%. Similar to Mn, Cr is an element effective for obtaining sufficient strength even at low C, and segregation is relatively small. Therefore, Cr is an element effective for ensuring strength and improving HIC resistance. If it is less than 0.05%, the effect is not sufficient, and if it exceeds 1%, weldability deteriorates, so the Cr content is specified to be 0.05 to 1%.
Ａｌ
Ａｌは、０．０７％以下とする。Ａｌは脱酸剤として添加されるが、０．０７％を超えると鋼の清浄度が低下し、耐ＨＩＣ性を劣化させるため、Ａｌ含有量は０．０７％以下に規定する。好ましくは、０．０１〜０．０７％とする。
Al
Al is made 0.07% or less. Al is added as a deoxidizer, but if it exceeds 0.07%, the cleanliness of the steel is lowered and the HIC resistance is deteriorated, so the Al content is specified to be 0.07% or less. Preferably, it is 0.01 to 0.07%.
Ｔｉ
Ｔｉは、０．０１〜０．０５％とする。Ｔｉは本発明において重要な元素で、Ｍｏと複合添加する。０．０１％以上添加することで、Ｍｏと複合析出物を形成し、強度上昇に大きく寄与する。
Ti
Ti is set to 0.01 to 0.05%. Ti is an important element in the present invention, and is added in combination with Mo. By adding 0.01% or more, a composite precipitate is formed with Mo, which greatly contributes to an increase in strength.
しかし、０．０５％を超えて添加すると、溶接熱影響部靭性の劣化を招くため、Ｔｉ含有量は０．０１〜０．０５％に規定する。さらに、０．０２５％未満にするとより優れた靭性を示す。このため、Ｎｂ、Ｖの一種または二種を添加する場合は、Ｔｉ含有量を０．０１〜０．０２５％未満とすることが好ましい。 However, if added over 0.05%, the weld heat affected zone toughness is deteriorated, so the Ti content is specified to be 0.01 to 0.05%. Furthermore, if it is less than 0.025%, more excellent toughness is exhibited. For this reason, when adding 1 type or 2 types of Nb and V, it is preferable to make Ti content into 0.01 to less than 0.025%.
Ｎ
Ｎは、０．００７％以下とする。０．００７％を超えて添加すると、溶接部の靭性劣化を招くと共に製鋼段階でのスラブ割れも招くため、含有量を０．００７％以下とする。
N
N is set to 0.007% or less. If added over 0.007%, the toughness of the welded part is deteriorated and slab cracking in the steel making stage is also caused, so the content is made 0.007% or less.
また、ＮはＴｉと共に析出物を形成するが、ＴｉＮ析出物は１３５０℃以上に達する溶接熱影響部の高温域において微細分散し、溶接熱影響部の旧オーステナイト粒を細粒化して、溶接熱影響部の靭性向上に大きく寄与する。その効果が十分となる含有量として０．００４〜０．００６％とすることが好ましい。 N forms precipitates with Ti, but the TiN precipitates are finely dispersed in the high temperature region of the weld heat affected zone reaching 1350 ° C. or more, and the old austenite grains in the weld heat affected zone are refined to produce welding heat. Significantly contributes to improved toughness of the affected area It is preferable to set it as 0.004 to 0.006% as content with which the effect becomes sufficient.
０．３５≦Ｃｅｑ１≦０．４５
但し、Ｃｅｑ１＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５とし
各元素は含有量（質量％）で含有しない元素は０とする。
0.35 ≦ Ceq1 ≦ 0.45
However, Ceq1 = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5, and each element is content (mass%), and elements not contained are 0.
ＡＰＩ Ｘ８０グレード等のＸ７０グレードを超える強度を確保するため、Ｃｅｑ１を０．３５以上とする。また、溶接性の観点から、良好な溶接性を確保するためにＣｅｑ１を０．４５以下とする。本発明に係る鋼材については、板厚１０ｍｍから３０ｍｍ程度の範囲でＣｅｑ１の板厚依存性はなく、３０ｍｍ程度まで同じＣｅｑ１で設計することができる。 In order to ensure strength exceeding X70 grade such as API X80 grade, Ceq1 is set to 0.35 or more. Further, from the viewpoint of weldability, Ceq1 is set to 0.45 or less in order to ensure good weldability. The steel material according to the present invention is not dependent on the plate thickness of Ceq1 in the range of plate thickness from 10 mm to 30 mm, and can be designed with the same Ceq1 up to about 30 mm.
Ｃｅｑ２≦０．６
但し、Ｃｅｑ２＝１．５Ｃ＋１．６Ｍｎ／６＋（１．３Ｃｕ＋１．３Ｎｉ）／１５＋（１．１Ｃｒ＋１．２Ｍｏ＋Ｖ）／５とし、各元素は含有量（質量％）で含有しない元素は０とする。
Ceq2 ≦ 0.6
However, Ceq2 = 1.5C + 1.6Mn / 6 + (1.3Cu + 1.3Ni) / 15 + (1.1Cr + 1.2Mo + V) / 5, and each element is content (mass%) and elements not contained are 0.
Ｃｅｑ２は合金元素の偏析度合いを示す指標であり、各合金元素の偏析係数を用いて下記式で表される。式において＊は乗算を示す。
Ｃｅｑ２＝ｆＣ＊Ｃ＋ ｆＭｎ＊Ｍｎ／６＋（ ｆＣｕ＊Ｃｕ＋ ｆＮｉ＊Ｎｉ）／１５＋（ ｆＣｒ＊Ｃｒ＋ ｆＭｏ＊Ｍｏ＋ ｆＶ＊Ｖ）／５
各合金元素の偏析係数はｆＣ＝１．５、ｆＭｎ＝１．６、ｆＣｕ＝１．３、ｆＮｉ＝１．３、ｆＣｒ＝１．１、ｆＭｏ＝１．２、ｆＶ＝１である。
Ceq2 is an index indicating the degree of segregation of the alloy elements, and is expressed by the following formula using the segregation coefficient of each alloy element. In the formula, * indicates multiplication.
Ceq2 = f C * C + f Mn * Mn / 6 + (f Cu * Cu + f Ni * Ni) / 15 + (f Cr * Cr + f Mo * Mo + f V * V) / 5
The segregation coefficient of each alloy element is f C = 1.5, f Mn = 1.6, f Cu = 1.3, f Ni = 1.3, f Cr = 1.1, f Mo = 1.2, f V = 1.
Ｃｅｑ２が０．６を超えると偏析部で硬化組織となり易く、割れ感受性が高まり耐ＨＩＣ特性が劣化するため、Ｃｅｑ２≦０．６とする。 If Ceq2 exceeds 0.6, a hardened structure is likely to be formed at the segregated portion, the cracking sensitivity is increased, and the HIC resistance is deteriorated, so Ceq2 ≦ 0.6.
Ｔｉ／Ｎ＞４
質量％でＴｉ量とＮ量の比：Ｔｉ／Ｎ＞４とする。本発明による高強度化はＴｉ、Ｍｏを含む複合析出物（主に炭化物）の微細析出によるものであるが、Ｔｉ／Ｎ≦４ではＴｉがＴｉＮの析出に消費されてしまい、析出強化に有効なＴｉ、Ｍｏを含む複合析出物を十分に得ることができないため、質量％でＴｉ量とＮ量の比：Ｔｉ／Ｎ＞４とする。
Ti / N> 4
Ratio of Ti amount and N amount in mass%: Ti / N> 4. Strengthening according to the present invention is due to fine precipitation of composite precipitates (mainly carbides) containing Ti and Mo. However, when Ti / N ≦ 4, Ti is consumed for precipitation of TiN, which is effective for precipitation strengthening. Since a composite precipitate containing Ti and Mo cannot be obtained sufficiently, the ratio of Ti amount to N amount by mass%: Ti / N> 4.
Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）
但し、各元素は原子％で含有しない元素は０とする。
C / (Mo + Ti + Nb + V)
However, the element which does not contain each element in atomic% is set to 0.
原子％でＣ量とＭｏ、Ｔｉ、Ｎｂ、Ｖの合計量の比である、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）を０．５〜３とする。本発明による高強度化はＴｉ、Ｍｏ，Ｎｂ，Ｖを含む複合析出物（主に炭化物）によるものである。Ｎｂ、Ｖの一種または二種を含む場合は後述する。 C / (Mo + Ti + Nb + V), which is the ratio of the amount of C and the total amount of Mo, Ti, Nb, and V in atomic%, is set to 0.5-3. The increase in strength according to the present invention is due to composite precipitates (mainly carbides) containing Ti, Mo, Nb, and V. The case of including one or two of Nb and V will be described later.
この複合析出物による析出強化を有効に利用するためには、Ｃ量と炭化物形成元素であるＭｏ、Ｔｉ、Ｎｂ，Ｖ量との関係が重要であり、これらの元素を適正なバランスのもとで添加することによって、熱的に安定かつ非常に微細な複合析出物を得ることが出来る。 In order to make effective use of precipitation strengthening by this composite precipitate, the relationship between the amount of C and the amounts of Mo, Ti, Nb, and V that are carbide forming elements is important. In this case, a thermally stable and very fine composite precipitate can be obtained.
各元素の原子％の含有量で表される、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）の値が０．５未満または３．０を超える場合はいずれかの元素量が過剰であり、硬化組織の形成による耐ＨＩＣ特性の劣化や靭性の劣化を招くため、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）の値を０．５〜３に規定する。 When the value of C / (Mo + Ti + Nb + V) represented by the atomic% content of each element is less than 0.5 or more than 3.0, the amount of any element is excessive, and the HIC resistance due to the formation of a hardened structure In order to cause deterioration of characteristics and toughness, the value of C / (Mo + Ti + Nb + V) is regulated to 0.5-3.
尚、質量％の含有量を用いる場合には（Ｃ／１２．０１）／（Ｍｏ／９５．９＋Ｔｉ／４７．９）の値を０．５〜３に規定する。Ｃ／（Ｍｏ＋Ｔｉ）の値を０．７〜２とすると、粒径５ｎｍ以下のより微細な析出物が得られるためより好ましい。 In addition, when using content of the mass%, the value of (C / 12.01) / (Mo / 95.9 + Ti / 47.9) is prescribed | regulated to 0.5-3. When the value of C / (Mo + Ti) is 0.7-2, a finer precipitate having a particle size of 5 nm or less is obtained, which is more preferable.
以上が本発明の基本成分組成であるが、本発明では鋼板の強度及び溶接部靭性を更に改善する目的で、Ｎｂ、Ｖの１種又は２種を含有してもよい。 The above is the basic component composition of the present invention. In the present invention, Nb and V may be contained in one or two kinds for the purpose of further improving the strength and weld zone toughness of the steel sheet.
Ｎｂ、Ｖの１種又は２種
Ｎｂ、Ｖの１種又は２種を含有する場合、Ｎｂは０．００５〜０．０７％とする。Ｎｂは組織の微細粒化により靭性を向上させるが、Ｔｉ及びＭｏと共に複合析出物を形成し強度上昇に寄与する。しかし、０．００５％未満では効果がなく、０．０７％を超えると溶接熱影響部の靭性が劣化するため、Ｎｂ含有量は０．００５〜０．０７％に規定する。
When 1 type or 2 types of Nb and V contain 1 type or 2 types of Nb and V, Nb shall be 0.005-0.07%. Nb improves toughness by refining the structure, but forms a composite precipitate with Ti and Mo and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.07%, the toughness of the weld heat affected zone deteriorates, so the Nb content is specified to be 0.005 to 0.07%.
Ｖは０．００５〜０．１％とする。Ｖは同様にＴｉおよびＭｏと共に複合析出物を形成し強度上昇に寄与する。しかし、０．００５％未満では効果がなく、０．１％を超えると溶接熱影響部の靭性が劣化するため、Ｖ含有量は０．００５〜０．１％に規定する。 V is 0.005 to 0.1%. V similarly forms a composite precipitate with Ti and Mo and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.1%, the toughness of the weld heat affected zone deteriorates, so the V content is specified to be 0.005 to 0.1%.
Ｎｂ、Ｖの一種または二種を含有する場合には、Ｃ量とＭｏ、Ｔｉ、Ｎｂ、Ｖの合計量の比である、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）を０．５〜３とする。 When one or two of Nb and V are contained, C / (Mo + Ti + Nb + V), which is a ratio of the amount of C and the total amount of Mo, Ti, Nb, and V, is set to 0.5 to 3.
本発明による高強度化はＴｉ、Ｍｏを含む析出物によるが、Ｎｂ、Ｖの一種または二種を含有する場合はそれらを含んだ複合析出物（主に炭化物）となる。このとき各元素の原子％の含有量で表される、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）の値が０．５未満または３を超える場合はいずれかの元素量が過剰であり、硬化組織の形成による耐ＨＩＣ特性の劣化や靭性の劣化を招くため、Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）の値を０．５〜３に規定する。 Strengthening according to the present invention depends on precipitates containing Ti and Mo, but when one or two of Nb and V are contained, they become composite precipitates (mainly carbides) containing them. At this time, when the value of C / (Mo + Ti + Nb + V) represented by the content of atomic% of each element is less than 0.5 or more than 3, the amount of any element is excessive, and HIC resistance due to formation of a hardened structure In order to cause deterioration of characteristics and toughness, the value of C / (Mo + Ti + Nb + V) is regulated to 0.5-3.
但し、各元素記号は原子％での含有量である。なお、質量％の含有量を用いる場合には（Ｃ／１２．０１）／（Ｍｏ／９５．９＋Ｔｉ／４７．９＋Ｎｂ／９２．９１＋Ｖ／５０．９４）の値を０．５〜３に規定する。より好ましくは、０．７〜２であり、粒径５ｎｍ以下のさらに微細な析出物が得られる。 However, each element symbol is a content in atomic%. In addition, when content of mass% is used, the value of (C / 12.01) / (Mo / 95.9 + Ti / 47.9 + Nb / 92.91 + V / 50.94) is specified to 0.5-3. . More preferably, it is 0.7-2, and a finer precipitate having a particle size of 5 nm or less is obtained.
本発明では鋼板の強度や耐ＨＩＣ特性をさらに改善する目的で、以下に示すＣｕ、Ｎｉ、Ｃａの１種または２種以上を含有してもよい。 In the present invention, for the purpose of further improving the strength and HIC resistance of the steel sheet, one or more of Cu, Ni and Ca shown below may be contained.
Ｃｕ
Ｃｕは、靭性の改善と強度の上昇に有効な元素であるが、多く添加すると溶接性が劣化するため、添加する場合は０．５％を上限とする。
Cu
Cu is an element effective for improving toughness and increasing strength, but if added in a large amount, weldability deteriorates, so when added, the upper limit is 0.5%.
Ｎｉ
Ｎｉは，靭性の改善と強度の上昇に有効な元素であるが、多く添加すると耐ＨＩＣ特性が低下するため、添加する場合は０．５％を上限とする。
Ni
Ni is an element effective for improving toughness and increasing strength, but if added in a large amount, the HIC resistance decreases, so when added, the upper limit is 0.5%.
Ｃａ
Ｃａは、硫化物系介在物の形態制御による耐ＨＩＣ特性向上に有効な元素であるが、０．０００５％未満ではその効果が十分でなく、０．００５％を超えて添加しても効果が飽和し、むしろ、鋼の清浄度の低下により耐ＨＩＣ性を劣化させるので、添加する場合はＣａ含有量を０．０００５〜０．００５％に規定する。
［製造条件］
図１に、本発明の組織制御方法を概略的に示す。上述した成分組成の鋼をＡｒ３以上のオーステナイト領域からベイナイト領域まで加速冷却することで、未変態オーステナイトとベイナイトの混合組織とする。
Ca
Ca is an element effective for improving the HIC resistance by controlling the form of sulfide inclusions. However, if it is less than 0.0005%, its effect is not sufficient, and even if added over 0.005%, it is effective. Saturated, but rather deteriorates the HIC resistance due to a decrease in the cleanliness of the steel, so when added, the Ca content is specified to be 0.0005 to 0.005%.
[Production conditions]
FIG. 1 schematically shows the tissue control method of the present invention. The steel having the above-described component composition is accelerated and cooled from the austenite region of Ar 3 or higher to the bainite region to obtain a mixed structure of untransformed austenite and bainite.
冷却後、直ちに再加熱する。未変態オーステナイトはフェライトに変態し、フェライト相中には微細析出物が分散析出する。一方、ベイナイト相は焼戻されて焼戻しベイナイトとなる。 Reheat immediately after cooling. Untransformed austenite is transformed into ferrite, and fine precipitates are dispersed and precipitated in the ferrite phase. On the other hand, the bainite phase is tempered to become tempered bainite.
微細析出物によって析出強化したフェライト相と焼戻されて軟化したベイナイト相の２相組織とすることで、高強度化と耐ＨＩＣ特性の両立が可能となる。 By adopting a two-phase structure of a ferrite phase that has been precipitation strengthened by fine precipitates and a bainite phase that has been tempered and softened, both high strength and HIC resistance can be achieved.
本発明に係るラインパイプ用高強度鋼板は上記の成分組成を有する鋼を用い、加熱温度：１０００〜１３００℃、圧延終了温度：Ａｒ３温度以上で熱間圧延を行い、その後５℃／ｓ以上の冷却速度で４００〜６００℃まで冷却し、冷却後直ちに０．５℃／ｓ以上の昇温速度で６００〜７００℃の温度まで再加熱を行うことで、ＭｏとＴｉを主体とする微細な複合炭化物を分散析出させ、ベイナイト相を軟化させた複合組織を備える。温度は鋼板の平均温度とする。 The high-strength steel sheet for line pipes according to the present invention uses steel having the above composition, and is hot-rolled at a heating temperature of 1000 to 1300 ° C. and a rolling end temperature of Ar 3 temperature or higher, and then 5 ° C./s or higher. By cooling to 400 to 600 ° C. at a cooling rate of 5 ° C., and immediately after cooling, it is reheated to a temperature of 600 to 700 ° C. at a temperature rising rate of 0.5 ° C./s or more. A composite structure in which composite carbide is dispersed and precipitated and the bainite phase is softened is provided. The temperature is the average temperature of the steel sheet.
加熱温度
加熱温度が１０００℃未満では炭化物の固溶が不十分で必要な強度が得られず、１３００℃を超えると靭性が劣化するため、１０００〜１３００℃とする。好ましくは、１０５０〜１２５０℃である。
Heating temperature If the heating temperature is less than 1000 ° C, the required strength cannot be obtained because the solid solution of the carbide is insufficient, and if it exceeds 1300 ° C, the toughness deteriorates. Preferably, it is 1050-1250 degreeC.
圧延終了温度
圧延終了温度がＡｒ３温度以下になると、その後のフェライト変態速度が低下するため、再加熱によるフェライト変態時に十分な微細析出物の分散析出が得られず、強度が低下するため、圧延終了温度をＡｒ３温度以上とする。
When the finish rolling temperature finish rolling temperature is below Ar 3 temperature, because the subsequent ferrite transformation rate decreases, dispersed precipitation of sufficient fine precipitates during the ferrite transformation caused by the reheating can not be obtained, because the strength is reduced, rolling The end temperature is set to Ar 3 temperature or higher.
Ａｒ３温度は、以下の式で求めることができる。
Ａｒ３＝９１０−３１０Ｃ−８０Ｍｎ−２０Ｃｕ−１５Ｃｒ−５５Ｎｉ−８０Ｍｏ
冷却条件
圧延終了後、直ちに５℃／ｓ以上の冷却速度で冷却し、４００〜６００℃で冷却停止する。圧延終了後に放冷または徐冷を行うと高温域から析出物が析出し、析出物が容易に粗大化し十分な強化が得られない。よって、析出強化に最適な温度まで急冷（加速冷却）を行い、高温域からの析出を防止する。
The Ar 3 temperature can be obtained by the following equation.
Ar 3 = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo
Cooling condition Immediately after the rolling, cooling is performed at a cooling rate of 5 ° C./s or more, and cooling is stopped at 400 to 600 ° C. When the product is allowed to cool or gradually cool after the completion of rolling, precipitates are precipitated from a high temperature range, and the precipitates are easily coarsened and sufficient strengthening cannot be obtained. Therefore, rapid cooling (accelerated cooling) is performed to a temperature optimal for precipitation strengthening, and precipitation from a high temperature region is prevented.
冷却速度が５℃／ｓ未満では高温域での析出防止効果が十分ではなく強度が低下するため、圧延終了後の冷却速度を５℃／ｓ以上に規定する。冷却方法は任意の冷却設備を用いることが可能である。 If the cooling rate is less than 5 ° C./s, the effect of preventing precipitation in a high temperature region is not sufficient and the strength is lowered. Therefore, the cooling rate after completion of rolling is specified to be 5 ° C./s or more. Any cooling equipment can be used as the cooling method.
冷却停止温度は４００〜６００℃とする。圧延終了後加速冷却し、ベイナイト変態域の４００〜６００℃まで急冷することにより、ベイナイト相を生成させ、かつ、再加熱時のフェライト変態の駆動力を大きくする。 The cooling stop temperature is 400 to 600 ° C. Accelerated cooling after the end of rolling and rapid cooling to 400 to 600 ° C. in the bainite transformation region generate a bainite phase and increase the driving force for ferrite transformation during reheating.
駆動力が大きくなることで、再加熱過程でのフェライト変態が促進され、短時間の再加熱でフェライト変態を完了させることが可能となる。冷却停止温度が４００℃未満では、ベイナイトかマルテンサイト単相組織、またはフェライト＋ベイナイト２相組織が得られても島状マルテンサイト（ＭＡ）が生成するために耐ＨＩＣ特性が劣化する。 By increasing the driving force, the ferrite transformation in the reheating process is promoted, and the ferrite transformation can be completed by reheating for a short time. When the cooling stop temperature is less than 400 ° C., even if a bainite, martensite single-phase structure, or a ferrite + bainite two-phase structure is obtained, island-like martensite (MA) is generated, so that the HIC resistance is deteriorated.
一方、６００℃を超えると再加熱時のフェライト変態が完了せずパーライトが析出し耐ＨＩＣ特性が劣化するため、加速冷却停止温度は４００〜６００℃に規定する。 On the other hand, if the temperature exceeds 600 ° C., ferrite transformation at the time of reheating is not completed and pearlite is precipitated, and the HIC resistance is deteriorated.
再加熱
加速冷却後直ちに０．５℃／ｓ以上の昇温速度で６００〜７００℃の温度まで再加熱を行う。本プロセスは本発明における重要な製造条件で、フェライト相の強化に寄与する微細析出物が、再加熱時のフェライト変態と同時に析出する。
Reheating Immediately after accelerated cooling, reheating is performed to a temperature of 600 to 700 ° C. at a heating rate of 0.5 ° C./s or more. This process is an important production condition in the present invention, and fine precipitates that contribute to strengthening of the ferrite phase are precipitated simultaneously with the ferrite transformation during reheating.
微細析出物によるフェライト相の強化とベイナイト相の軟化を同時に行い、フェライト相とベイナイト相の強度差の小さい組織を得るためには、加速冷却後直ちに６００〜７００℃の温度域まで再加熱することが必要である。 In order to simultaneously strengthen the ferrite phase with fine precipitates and soften the bainite phase and obtain a structure with a small difference in strength between the ferrite phase and the bainite phase, reheat to 600 to 700 ° C immediately after accelerated cooling. is required.
また、再加熱は、冷却後の温度より少なくとも５０℃以上昇温することが望ましい。再加熱時の昇温速度が０．５℃／ｓ未満では、目的の再加熱温度に達するまでに長時間を要するため製造効率が悪化し、またパーライト変態が生じるため、微細析出物の分散析出が得られず十分な強度を得る事ができない。 Further, it is desirable that the reheating is performed at least 50 ° C. higher than the temperature after cooling. When the heating rate during reheating is less than 0.5 ° C./s, it takes a long time to reach the target reheating temperature, so that production efficiency deteriorates and pearlite transformation occurs, so that fine precipitates are dispersed and precipitated. Cannot be obtained and sufficient strength cannot be obtained.
再加熱温度が６００℃未満ではフェライト変態が完了せずその後の冷却時に未変態オーステナイトがパーライトに変態するため耐ＨＩＣ特性が劣化する。 If the reheating temperature is less than 600 ° C., the ferrite transformation is not completed, and untransformed austenite transforms to pearlite during subsequent cooling, so that the HIC resistance is deteriorated.
一方、７００℃を超えると析出物が粗大化し十分な強度が得られないため、再加熱温度域を６００〜７００℃に規定する。 On the other hand, if the temperature exceeds 700 ° C., the precipitates become coarse and sufficient strength cannot be obtained. Therefore, the reheating temperature range is regulated to 600 to 700 ° C.
本発明の製造方法を用いれば再加熱後直ちに冷却しても、フェライト変態が十分に進行するため、再加熱温度における保持時間は規定しない。 Even if it cools immediately after reheating if the manufacturing method of this invention is used, since ferrite transformation will fully advance, the holding time in reheating temperature is not prescribed | regulated.
但し、３０分を超えて温度保持を行うと、析出物の粗大化を生じ強度低下を招く場合があるため、３０分以内とすることが好ましい。 However, if the temperature is maintained for more than 30 minutes, precipitates may be coarsened and the strength may be reduced.
再加熱後の冷却速度は適宜設定すれば良く規定しない。但し、再加熱後の冷却過程でもフェライト変態が進行するので、空冷が好ましい。フェライト変態を阻害しない程度であれば、空冷よりも早い冷却速度で冷却を行うことも可能である。 The cooling rate after reheating is not specified as long as it is set appropriately. However, air cooling is preferable because the ferrite transformation proceeds in the cooling process after reheating. As long as the ferrite transformation is not hindered, it is possible to perform cooling at a cooling rate faster than air cooling.
６００〜７００℃の温度まで再加熱を行うための設備として、加速冷却設備の下流側に加熱装置を設置することができる。加熱装置としては、鋼板の急速加熱が可能である誘導加熱装置やガス燃焼炉を用いる事が好ましい。 As equipment for performing reheating to a temperature of 600 to 700 ° C., a heating device can be installed on the downstream side of the accelerated cooling equipment. As a heating device, it is preferable to use an induction heating device or a gas combustion furnace capable of rapid heating of a steel plate.
誘導加熱装置は均熱炉等に比べて温度制御が容易でありコストも比較的低く、冷却後の鋼板を迅速に加熱できるので特に好ましい。また複数の誘導加熱装置を直列に連続して配置することにより、ライン速度や鋼板の種類・寸法が異なる場合にも、通電する誘導加熱装置の数や供給電力を任意に設定するだけで、昇温速度、再加熱温度を自在に操作することが可能である。なお、再加熱後の冷却速度は任意の速度で構わないので、加熱装置の下流側には特別な設備を設置する必要はない。 The induction heating device is particularly preferable because temperature control is easier than in a soaking furnace, the cost is relatively low, and the cooled steel sheet can be heated quickly. In addition, by arranging a plurality of induction heating devices in series, even if the line speed and the type and size of the steel sheet are different, the number of induction heating devices to be energized and the supply power can be set by arbitrarily setting them. It is possible to freely control the temperature rate and the reheating temperature. In addition, since the cooling rate after reheating may be arbitrary, it is not necessary to install special equipment in the downstream of a heating apparatus.
図２に、本発明の製造方法を実施するための製造ラインの一例の概略図を示す。圧延ライン１には上流から下流側に向かって熱間圧延機３、加速冷却装置４、インライン型誘導加熱装置５、ホットレベラー６が配置されている。
In FIG. 2, the schematic of an example of the manufacturing line for enforcing the manufacturing method of this invention is shown. In the rolling line 1, a
インライン型誘導加熱装置５あるいは他の熱処理装置を、圧延設備である熱間圧延機３およびそれに引き続く冷却設備である加速冷却装置４と同一ライン上に設置する事によって、圧延、冷却終了後迅速に再加熱処理が行えるので、圧延して加速冷却した後の鋼板を、直ちに６００℃以上に加熱することが可能である。
By installing the in-line type
本発明に係る鋼板は、従来の加速冷却等で得られるベイナイトまたはアシキュラーフェライト組織の鋼板のような表層部での硬さの上昇がないので、表層部からのＨＩＣが生じない。 Since the steel sheet according to the present invention does not increase in hardness at the surface layer portion unlike a conventional bainite or acicular ferrite structure steel plate obtained by accelerated cooling or the like, HIC from the surface layer portion does not occur.
更に強度差の小さいフェライト相とベイナイト相の２相組織は割れに対する抵抗が極めて高く、ＨＩＣ発生の起点や伝播経路と成り得る硬質相のＭＡの生成を低減できるため、鋼板中心部や介在物からのＨＩＣも抑制することが可能となる。 Furthermore, the two-phase structure of the ferrite phase and bainite phase, which have a small strength difference, has extremely high resistance to cracking and can reduce the generation of hard phase MA that can serve as the origin and propagation path of HIC generation. HIC can also be suppressed.
そのため、プレスベンド成形、ロール成形、ＵＯＥ成形等で鋼管に成形して、原油や天然ガスを輸送する鋼管（電縫鋼管、スパイラル鋼管、ＵＯＥ鋼管）として好適である。 Therefore, it is suitable as a steel pipe (an electric resistance steel pipe, a spiral steel pipe, a UOE steel pipe) that is formed into a steel pipe by press bend molding, roll molding, UOE molding, or the like and transports crude oil or natural gas.
表１に示す化学成分の鋼（鋼種Ａ〜Ｎ）を連続鋳造法によりスラブとし、板厚２２、２
８ｍｍの厚鋼板（Ｎｏ．１〜２４（Ｎｏ．２は欠番））を製造した。
Steel of the chemical composition shown in Table 1 (steel types A to N) was made into a slab by a continuous casting method, and plate thicknesses 22 and 2
8 mm thick steel plates ( No. 1 to 24 (No. 2 is a missing number) ) were manufactured.
加熱したスラブを熱間圧延により圧延した後、直ちに水冷型の加速冷却設備を用いて冷却を行い、誘導加熱炉またはガス燃焼炉を用いて再加熱を行った。冷却設備及び誘導加熱炉はインライン型とした。 After the heated slab was rolled by hot rolling, it was immediately cooled using a water-cooled accelerated cooling facility and reheated using an induction heating furnace or a gas combustion furnace. The cooling equipment and induction heating furnace were in-line type.
以上のようにして製造した鋼板のミクロ組織を、光学顕微鏡（倍率×４００）、透過型電子顕微鏡（ＴＥＭ）により観察した。観察は圧延方向と直角方向となる断面について板厚１／２で面積１２ｍｍ2について１０視野行い、各視野の平均値とした。 The microstructure of the steel sheet produced as described above was observed with an optical microscope (magnification × 400) and a transmission electron microscope (TEM). Observation was carried out for a cross section perpendicular to the rolling direction and 10 views for an area of 12 mm 2 with a plate thickness of ½, and the average value for each view.
析出物の成分はエネルギー分散型Ｘ線分光法（ＥＤＸ）により分析した。また各鋼板の引張特性、耐ＨＩＣ特性を測定した。 The components of the precipitate were analyzed by energy dispersive X-ray spectroscopy (EDX). The tensile properties and HIC resistance of each steel plate were measured.
引張特性は、圧延垂直方向の全厚試験片を引張試験片として引張試験を行い、降伏強度、引張強度を測定した。 Tensile properties were measured by performing a tensile test using a full thickness test piece in the rolling vertical direction as a tensile test piece, and measuring yield strength and tensile strength.
そして、製造上のばらつきを考慮して、降伏強度５００ＭＰａ以上、引張強度６００ＭＰａ以上であるものをＡＰＩ Ｘ７０グレード以上の高強度鋼板として評価した。 In consideration of manufacturing variations, a steel having a yield strength of 500 MPa or more and a tensile strength of 600 MPa or more was evaluated as a high strength steel plate of API X70 grade or more.
耐ＨＩＣ特性はＮＡＣＥ Ｓｔａｎｄａｒｄ ＴＭ−０２−８４に準じた浸漬時間９６時間のＨＩＣ試験を行い、割れが認められない場合を耐ＨＩＣ性良好と判断して○で、割れが発生した場合を×で示した。 The HIC resistance was evaluated by performing an HIC test with an immersion time of 96 hours in accordance with NACE Standard TM-02-84. If no cracks were observed, the HIC resistance was judged good. Indicated.
表２に、各鋼板（Ｎｏ．１〜２４（Ｎｏ．２は欠番））の製造条件と測定結果を併せて示す。本発明例であるＮｏ．１〜１３（Ｎｏ．２は欠番）はいずれも、化学成分および製造方法が本発明の範囲内であり、降伏強度５００ＭＰａ以上、引張強度６００ＭＰａ以上の高強度で、かつ耐ＨＩＣ性が優れていた。 In Table 2, manufacturing conditions and measurement results of each steel plate ( No. 1 to 24 (No. 2 is a missing number) ) are shown together. No. which is an example of the present invention. 1 to 13 (No. 2 is a missing number) all have chemical components and production methods within the scope of the present invention, yield strength of 500 MPa or higher, tensile strength of 600 MPa or higher, and excellent HIC resistance. .
ＴｉとＭｏと、一部の鋼板についてはさらにＮｂおよび／またはＶを含む粒径が１０ｎｍ未満の微細な炭化物の析出物が分散析出していた。また、鋼板の組織は、実質的にフェライト＋ベイナイト２相組織であり、ベイナイト相の分率はいずれも１０〜８０％の範囲であり、ＭＡの分率はいずれも５％以下であった。 As for Ti, Mo, and some of the steel plates, fine carbide precipitates containing Nb and / or V and having a particle size of less than 10 nm were dispersed and precipitated. Moreover, the structure of the steel sheet was substantially a ferrite + bainite two-phase structure, the fraction of the bainite phase was in the range of 10 to 80%, and the fraction of MA was 5% or less.
Ｎｏ．１４〜１８は、化学成分は本発明の範囲内であるが、製造方法が本発明の範囲外であるため、組織がフェライト＋ベイナイト２相組織になっていないことや、微細炭化物が分散析出していないため、強度不足やＨＩＣ試験で割れが発生した。 No. In Nos. 14 to 18, the chemical components are within the scope of the present invention, but the manufacturing method is outside the scope of the present invention, so that the structure is not a ferrite + bainite two-phase structure, and fine carbides are dispersed and precipitated. Therefore, cracks occurred in the strength and HIC tests.
Ｎｏ．１９〜２４は化学成分が本発明の範囲外であるので、粗大な析出物が生成したり、ＴｉとＭｏを含む析出物が分散析出していないため、十分な強度が得られないか、ＭＡの過剰生成によりＨＩＣ試験で割れが生じた。 No. Since the chemical components of Nos. 19 to 24 are outside the scope of the present invention, coarse precipitates are not formed, and precipitates containing Ti and Mo are not dispersed and precipitated. Cracking occurred in the HIC test due to excessive formation of.
なお、再加熱を誘導加熱炉で行った場合もガス燃焼炉で行った場合も特に結果に差は見られなかった。 In addition, when the reheating was performed in the induction heating furnace or in the gas combustion furnace, there was no particular difference in the results.
１：圧延ライン
２：鋼板
３：熱間圧延機
４：加速冷却装置
５：インライン型誘導加熱装置
６：ホットレベラー
1: Rolling line 2: Steel plate 3: Hot rolling mill 4: Accelerated cooling device 5: In-line induction heating device 6: Hot leveler
Claims (6)
０．５〜１．８％、Ｐ：０．０１％以下、Ｓ：０．００２％以下、Ｍｏ：０．０５〜０．
５％、Ｃｒ：０．０５〜１％、Ａｌ：０．０７％以下、Ｎｂ：０．００５〜０．０７％、Ｔｉ：０．０１〜０．０５％、Ｃａ：０．０００５〜０．００５％、Ｎ：０．００７％以下、下記（１）〜（４）式を満足し、残部Ｆｅおよび不可避的不純物で
金属組織がフェライトとベイナイトの２相組織であり、島状マルテンサイト（ＭＡ）の分
率が体積分率で５％以下であり、フェライト相に、ＴｉとＭｏを含む複合析出物が分散析出していることを
特徴とする、耐ＨＩＣ特性に優れたラインパイプ用高強度鋼板。
０．３５≦Ｃｅｑ１≦０．４５・・・（１）
但し、Ｃｅｑ１＝Ｃ＋Ｍｎ／６＋（Ｃｕ＋Ｎｉ）／１５＋（Ｃｒ＋Ｍｏ＋Ｖ）／５
で、各元素は含有量（質量％）で含有しない元素は０とする。
Ｃｅｑ２≦０．６・・・（２）
但し、Ｃｅｑ２＝１．５Ｃ＋１．６Ｍｎ／６＋（１．３Ｃｕ＋１．３Ｎｉ）／１５＋（１
．１Ｃｒ＋１．２Ｍｏ＋Ｖ）／５
但し、各元素は含有量（質量％）で含有しない元素は０とする。
Ｔｉ／Ｎ＞４・・・（３）
但し、Ｔｉ，Ｎは質量％とする。
０．５≦Ｃ／（Ｍｏ＋Ｔｉ＋Ｎｂ＋Ｖ）≦３・・・（４）
但し、各元素は原子％で含有しない元素は０とする。 Steel composition is mass%, C: 0.02-0.08%, Si: 0.01-0.5%, Mn:
0.5-1.8%, P: 0.01% or less, S: 0.002% or less, Mo: 0.05-0.
5%, Cr: 0.05-1%, Al: 0.07% or less, Nb: 0.005-0.07%, Ti: 0.01-0.05%, Ca: 0.0005-0. 005%, N: 0.007% or less, satisfying the following formulas (1) to (4), the balance Fe and unavoidable impurities, the metal structure is a two-phase structure of ferrite and bainite, and island martensite (MA ) With a volume fraction of 5% or less, and composite precipitates containing Ti and Mo are dispersed and precipitated in the ferrite phase . High strength for line pipes with excellent HIC resistance steel sheet.
0.35 ≦ Ceq1 ≦ 0.45 (1)
However, Ceq1 = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5
Thus, each element is contained in an amount (mass%), and elements not contained are set to 0.
Ceq2 ≦ 0.6 (2)
However, Ceq2 = 1.5C + 1.6Mn / 6 + (1.3Cu + 1.3Ni) / 15 + (1
. 1Cr + 1.2Mo + V) / 5
However, each element is contained (mass%), and elements not contained are set to 0.
Ti / N> 4 (3)
However, Ti and N are mass%.
0.5 ≦ C / (Mo + Ti + Nb + V) ≦ 3 (4)
However, the element which does not contain each element in atomic% is set to 0.
００〜１３００℃、圧延終了温度：Ａｒ３温度以上の条件で熱間圧延した後、冷却速度：
５℃／ｓ以上で４００〜６００℃まで加速冷却を行い、冷却後直ちに昇温速度：０．５℃
／ｓ以上で６００〜７００℃の温度まで再加熱を行うことを特徴とする、耐ＨＩＣ特性に
優れたラインパイプ用高強度鋼板の製造方法。 The steel containing the chemical component according to any one of claims 1 to 3 is heated at a heating temperature of 10
After hot rolling under conditions of 00 to 1300 ° C. and rolling end temperature: Ar 3 temperature or higher, cooling rate:
Accelerated cooling to 400 to 600 ° C. at 5 ° C./s or more, and immediately after cooling, the rate of temperature increase: 0.5 ° C.
The manufacturing method of the high strength steel plate for line pipes excellent in the HIC-proof characteristic characterized by performing reheating to the temperature of 600-700 degreeC above / s.
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Family Cites Families (5)
| Publication number | Priority date | Publication date | Assignee | Title |
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| JPH07216500A (en) * | 1994-01-28 | 1995-08-15 | Sumitomo Metal Ind Ltd | High strength steel material excellent in corrosion resistance and its production |
| JP4419695B2 (en) * | 2003-06-12 | 2010-02-24 | Ｊｆｅスチール株式会社 | Low yield ratio high strength high toughness steel sheet and method for producing the same |
| JP4507746B2 (en) * | 2003-07-31 | 2010-07-21 | Ｊｆｅスチール株式会社 | Low yield ratio high strength high toughness steel pipe excellent in strain aging resistance and method for producing the same |
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