JP5369639B2 - High strength steel material excellent in welding heat-affected zone toughness and HIC resistance and manufacturing method thereof - Google Patents
High strength steel material excellent in welding heat-affected zone toughness and HIC resistance and manufacturing method thereof Download PDFInfo
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本発明は、ラインパイプ用鋼管等に用いられる、製造・施工に溶接を伴う高強度鋼材に関するものである。 The present invention relates to a high-strength steel material that is used for a steel pipe for a line pipe and the like and that involves welding during manufacture and construction.
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には、強度、靭性、溶接性の他、耐水素誘起割れ性(以下、耐HIC性と称す)や耐応力腐食割れ性(以下、耐SCC性と称す)などの諸特性が必要とされる。特に、近年は、国際的な大規模ラインパイプ事業の展開に伴って、アメリカ石油協会(以下、APIと称す)規格 X80グレード等のX65グレードを超えるラインパイプ用高強度鋼板のニーズが高まりつつある。 In addition to strength, toughness and weldability, line pipes used to transport crude oil and natural gas containing hydrogen sulfide have resistance to hydrogen-induced cracking resistance (hereinafter referred to as HIC resistance) and stress corrosion cracking resistance (hereinafter referred to as Various characteristics such as SCC resistance) are required. In particular, in recent years, with the development of international large-scale line pipe business, the need for high-strength steel sheets for line pipes exceeding X65 grade such as American Petroleum Institute (hereinafter referred to as API) standard X80 grade is increasing. .
製造や施工に溶接を伴うこの種の鋼材を高強度化する場合、靭性、特に、溶接熱影響部(以下、HAZと称す)の靭性を劣化させないことが重要である。この強度とHAZ靭性の両立という制約条件があるため、従来、主に、固溶強化、結晶粒微細化、複相組織化によって要求強度を達成する材料設計がおこなわれており、析出強化は靭性を損なう可能性が高いので積極的には利用されていないのが現状である。 When increasing the strength of this type of steel that involves welding during manufacturing and construction, it is important not to deteriorate the toughness, particularly the toughness of the heat affected zone (hereinafter referred to as HAZ). Because there is a constraint condition that this strength and HAZ toughness are compatible, conventionally, material design that achieves the required strength mainly by solid solution strengthening, grain refinement, and multiphase structure has been performed, and precipitation strengthening is toughness. The current situation is that it is not actively used because it is likely to damage
一方、従来技術でHAZ靭性を損なうことなく、API規格X65グレード以上の強度を達成しようとすると高めの合金添加量が必要になって耐HIC特性が劣化するという課題がある。鋼材のHICは、腐食反応によって鋼材表面に吸着した水素が、原子状水素として鋼板内部に侵入し、鋼中のMnSなどの非金属介在物や硬質な第2相組織の周りに拡散・集積し、その内圧により割れを生ずるものとされている。HICを防ぐ方法としては、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方法が知られている(例えば、特許文献1)。耐HIC性の優れたX80グレードの高強度鋼板に関しては、低SでCa添加により介在物の形態制御を行いつつ、低C、低Mnとして中心偏析を抑制し、それに伴う強度低下をCr、Niなどの添加と加速冷却によって補う方法が知られている(特許文献2〜4)。また、ミクロ組織をフェライト単相組織とすることで耐SCC性や耐HIC性を改善し、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を利用した高強度鋼も開示されている(例えば、特許文献5)。金属組織が実質的にフェライトとベイナイトの2相組織で、TiとMoとを含む析出物が分散していることを特徴とする耐HIC性とHAZ靭性に優れたラインパイプ用高強度鋼板とその製造方法も開示されている(例えば、特許文献6および7)。
On the other hand, there is a problem in that the HIC resistance is deteriorated when a higher alloy addition amount is required to achieve the strength of API standard X65 grade or higher without impairing the HAZ toughness with the prior art. In HIC of steel, hydrogen adsorbed on the steel surface due to corrosion reaction penetrates into the steel sheet as atomic hydrogen, and diffuses and accumulates around non-metallic inclusions such as MnS and hard second phase structure in the steel. The internal pressure is considered to cause cracks. As a method to prevent HIC, by adding appropriate amount of Ca and Ce to the amount of S, the formation of acicular MnS is suppressed, and the shape is changed to finely dispersed spherical inclusions with small stress concentration and cracking There is known a method for suppressing the propagation of the occurrence of this (for example, Patent Document 1). For high strength steel of X80 grade with excellent HIC resistance, while controlling the form of inclusions by adding Ca with low S, it suppresses center segregation as low C, low Mn, and the accompanying strength reduction is Cr, Ni The method of supplementing by adding such as and accelerated cooling is known (
しかし、特許文献1〜4に記載された耐HIC性を改善する方法はいずれも中心偏析部が対象である。API X80グレード等のX65グレードを超える高強度鋼板は加速冷却または直接焼入れによって製造される場合が多いため、冷却速度の速い鋼板表面部が内部に比べ硬化し、表面近傍から水素誘起割れが発生する。また、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であり、中心偏析部のHICへの対策を施した場合でも、API X80グレード程度の高強度鋼では硫化物系または酸化物系介在物を起点としたHICをなくすことは困難である。 However, the methods for improving the HIC resistance described in Patent Documents 1 to 4 are all about the center segregation part. High strength steel sheets exceeding the X65 grade, such as API X80 grade, are often manufactured by accelerated cooling or direct quenching, so the surface of the steel sheet with a high cooling rate is hardened compared to the inside, and hydrogen-induced cracking occurs from the vicinity of the surface. . Moreover, 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, and measures against HIC at the center segregation part were taken. Even in this case, it is difficult for high strength steel of about API X80 grade to eliminate HIC starting from sulfide or oxide inclusions.
また、われわれの調査によると、特許文献2〜4の方法でX80グレードの強度を安定的に確保するためには、炭素当量Ceq(Ceq=C+Mn/6+Cu/15+Ni/15+Cr/5+Mo/5+V/5)を0.41以上に設定する必要がある。しかし、この炭素当量Ceqが0.41以上の設定値では、表層や円周溶接部の最高硬さが高くなってHAZ靭性を安定的に確保できなくなるため、この方法は鋼板の強度とHAZ靭性を両立できる技術とは考えられない。
According to our research, in order to stably secure the strength of X80 grade by the methods of
特許文献5に記載の高強度鋼で利用している微細なTiCは析出強化に有効だが、溶接時に溶解・再析出を経ると粗大化しやすいため、HAZ靭性を劣化させやすいという問題がある。 The fine TiC used in the high-strength steel described in Patent Document 5 is effective for precipitation strengthening, but it tends to coarsen after melting and reprecipitation at the time of welding, so that there is a problem that HAZ toughness is likely to deteriorate.
HAZ靭性の向上に関してTiに着目した特許文献6および7では、Tiが50〜250 mass ppmの範囲に規定されているが、Tiと優先的に窒化物を形成するNが40〜60mass ppmの範囲に規定されているから、添加したTiのうち、140〜200mass ppmはTiNとして析出すると予測できる。したがって、同文献に記載された技術では、TiNの微細分散による旧オーステナイトの細粒化を利用したHAZ靭性の向上は期待できるとしても、Ti、Mo、Nb、Vを含む複合炭化物を析出させられるだけの固溶Tiを安定して確保できにくいと予想される。
本発明の目的は、このような従来技術の課題を解決し、耐HIC特性を損なうことなく、API規格X80グレード以上の強度と溶接熱影響部の靭性とを兼ね備えたラインパイプ用鋼管等に用いられる高強度鋼材を提供することにある。 The object of the present invention is to solve such problems of the prior art and use it for steel pipes for line pipes, etc. that have both the strength of API standard X80 grade or higher and the toughness of weld heat affected zone without impairing HIC resistance. It is to provide a high-strength steel material that can be manufactured.
本発明者が、鋼材の強度とHAZ靭性を両立できる材料設計について鋭意研究を重ねた結果、HAZで粗大化した析出物の原因となるNbの含有量も、HAZ靭性を損なうCeqも増やすことなく、析出強化を効果的に実現する方法を独自に見い出した。 As a result of intensive research on material design that can achieve both steel strength and HAZ toughness, the present inventor has found that the Nb content that causes precipitates coarsened by HAZ and the Ceq that impairs HAZ toughness are not increased. In addition, we have found a unique method for effectively realizing precipitation strengthening.
具体的には、(1)スラブ加熱段階で固溶Tiを確保するように成分設計し、圧延後、再加熱処理をすると、Ti系析出物の析出に牽引されるかのように、含有したNbの析出物の量が格段に増加すること(以下、これを「Tiの析出牽引効果」と呼ぶ)、(2)その析出過程で形成される、TiよりもNbがリッチで大きさが20nm未満の微細複合析出物を一定量以上確保すれば、Nbの含有量も炭素当量Ceqも増やすことなく、強度を効果的に向上させられること、を見い出した。
本発明は、以上の知見に基づいてなされたものであり、その要旨は以下のとおりである。
[1]C: 0.02〜0.08 mass%、Si: 0.01〜0.5 mass%、Mn: 0.5〜2.0 mass%、Ca: 0.0005〜0.003mass%、Ti: 0.01〜0.03 mass%、Nb: 0.04〜0.05 mass%、Al:0.07mass%以下、Cu:0.5mass%以下、Ni:0.5mass%以下、Cr: 0.5mass%以下、N: 0.007 mass%以下、残部がFeおよび不可避的不純物からなる鋼材で、炭素当量Ceqが0.41未満、TiはNの3.4倍に0.003mass%を加えた値以上を含有し、かつ、大きさ20nm未満の析出物中のTiが鋼材全体に対して10mass ppm以上、Nbが鋼材全体に対して140mass ppm以上であることを特徴とする溶接熱影響部靭性と耐HIC特性に優れた高強度鋼材。
[2]前記[1]において、さらに、Mo: 0.05〜0.5 mass%、V: 0.005〜0.1mass%のいずれか1つ以上を含み、大きさ20nm未満の析出物中のMoが鋼材全体に対して50mass ppm以上および/またはVが鋼材全体に対して5mass ppm以上であることを特徴とする請求項1記載の溶接熱影響部靭性と耐HIC特性に優れた高強度鋼材。
[3]前記[1]または[2]に記載の化学成分を含有する鋼を、加熱温度:1050〜1250℃、圧延終了温度:Ar3 温度以上の温度域で熱間圧延した後、冷却速度:5℃/sec以上で300〜600℃まで加速冷却を行い、冷却停止温度で0.5〜3分間放冷し、次いで、昇温速度:0.5〜2.0℃/secで、最高到達温度:600〜700℃まで再加熱を行った後、ただちに空冷以上の冷却速度で冷却することを特徴とする溶接熱影響部靭性と耐HIC特性に優れた高強度鋼材の製造方法。
なお、本明細書において、鋼の成分を示す%、ppmは、すべてmass%、mass ppmである。また、Ar3 温度は冷却時のフェライト変態開始温度であり、Ar3=910-310C-80Mn-20Cu-15Cr-55Ni-80Mo から求めることができる。
Specifically, (1) component design to ensure solute Ti in the slab heating stage, and after rolling, reheat treatment, contained as if pulled by the precipitation of Ti-based precipitates The amount of Nb precipitates increases dramatically (hereinafter referred to as the “Ti precipitation traction effect”), and (2) Nb is richer than Ti and 20 nm in size during the precipitation process. It was found that the strength can be effectively improved without increasing the Nb content and the carbon equivalent Ceq if a certain amount or more of fine composite precipitates is secured.
This invention is made | formed based on the above knowledge, The summary is as follows.
[1] C: 0.02 to 0.08 mass%, Si: 0.01 to 0.5 mass%, Mn: 0.5 to 2.0 mass%, Ca: 0.0005 to 0.003 mass%, Ti: 0.01 to 0.03 mass%, Nb: 0.04 to 0.05 mass% , Al: 0.07 mass% or less, Cu: 0.5 mass% or less, Ni: 0.5 mass% or less, Cr: 0.5 mass% or less, N: 0.007 mass% or less, the balance being Fe and inevitable impurities steel material, carbon equivalent Ceq is less than 0.41, Ti contains 3.4 times N or more than 0.003 mass%, and Ti in precipitates less than 20 nm in size is 10 mass ppm or more, and Nb is the whole steel. High strength steel with excellent weld heat affected zone toughness and HIC resistance, characterized by being 140 mass ppm or higher.
[2] In the above [1], Mo in the precipitate further including any one or more of Mo: 0.05 to 0.5 mass%, V: 0.005 to 0.1 mass% is less than 20 nm in size relative to the whole steel material. The high-strength steel material excellent in weld heat-affected zone toughness and HIC resistance according to claim 1, characterized in that it is 50 mass ppm or more and / or V is 5 mass ppm or more with respect to the whole steel material.
[3] A steel containing the chemical component described in [1] or [2] above is hot-rolled at a heating temperature: 1050 to 1250 ° C. and a rolling end temperature: Ar 3 temperature or higher, and then a cooling rate: Accelerated cooling to 300-600 ° C at 5 ° C / sec or more, left to cool for 0.5-3 minutes at the cooling stop temperature, and then the highest temperature: 600-700 ° C at a heating rate of 0.5-2.0 ° C / sec A method for producing high-strength steel with excellent weld heat-affected zone toughness and HIC resistance, which is characterized by cooling immediately after reheating up to a cooling rate higher than air cooling.
In this specification, “%” and “ppm” indicating the components of steel are mass% and mass ppm, respectively. The Ar3 temperature is the ferrite transformation start temperature during cooling, and can be determined from Ar3 = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo.
本発明を適用すれば、耐HIC特性が劣化しない低合金成分で、API規格X65グレード以上の強度と溶接熱影響部の靭性を両立させることが格段に容易になる。 By applying the present invention, it is much easier to achieve both the strength of API standard X65 grade or higher and the toughness of the weld heat affected zone with a low alloy component that does not deteriorate the HIC resistance.
本発明は、スラブ加熱段階で固溶Tiを確保する成分設計と再加熱時のTiの析出牽引効果を利用して、TiよりもNbがリッチで大きさが20nm未満の微細析出物を一定量以上析出させると、HAZ靭性や耐HIC特性に懸念のあるNb含有量も炭素当量Ceqも増やすことなく、強度上昇を達成できるという点にある。 The present invention uses a component design that ensures solid solution Ti in the slab heating stage and the precipitation traction effect of Ti during reheating, so that a certain amount of fine precipitates with Nb richer than Ti and less than 20 nm in size are obtained. Precipitation as described above is that strength can be increased without increasing the Nb content and the carbon equivalent Ceq, which are concerned about HAZ toughness and HIC resistance.
1)まず、本発明で用いる高強度鋼材の化学成分について説明する。
C: 0.02〜0.08 mass%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、0.02mass%未満では十分な強度が確保できず、0.08mass%を超えると靭性や耐HIC性を劣化させるため、C含有量はこの範囲が好ましい。更に好ましくは、0.03〜0.06mass%である。
1) First, chemical components of the high-strength steel material used in the present invention will be described.
C: Set to 0.02 to 0.08 mass%.
C is an element that contributes to precipitation strengthening as a carbide, but if it is less than 0.02 mass%, sufficient strength cannot be secured, and if it exceeds 0.08 mass%, the toughness and HIC resistance deteriorate, so the C content is within this range. preferable. More preferably, it is 0.03-0.06 mass%.
Si: 0.01〜0.5 mass%とする。
Siは脱酸のために含有するが、0.01mass%未満では脱酸効果が十分でなく、0.5mass%を超えると靭性や溶接性を劣化させるため、Si含有量はこの範囲が好ましい。更に好ましくは、0.01〜0.3mass%である。
Si: 0.01 to 0.5 mass%.
Si is contained for deoxidation, but if it is less than 0.01 mass%, the deoxidation effect is not sufficient, and if it exceeds 0.5 mass%, the toughness and weldability are deteriorated, so the Si content is preferably in this range. More preferably, it is 0.01-0.3 mass%.
Mn: 0.5〜2.0 mass%とする。
Mnは強度、靭性のために含有するが、0.5mass%未満ではその効果が十分ではなく、2.0mass%を超えると溶接性と耐HIC性が劣化するため、Mn含有量はこの範囲が好ましい。更に好ましくは、0.5〜1.5mass%である。
Mn: 0.5 to 2.0 mass%.
Mn is contained for strength and toughness, but if it is less than 0.5 mass%, the effect is not sufficient, and if it exceeds 2.0 mass%, the weldability and HIC resistance deteriorate, so the Mn content is preferably within this range. More preferably, it is 0.5 to 1.5 mass%.
Al: 0.07mass%以下とする。
Alは脱酸剤として含有するが、0.07mass%を超えると鋼の清浄度が低下し、耐HIC特性を劣化させるため、Al含有量はこの範囲が好ましい。更に好ましくは、0.01〜0.07mass%である。
Al: 0.07 mass% or less.
Al is contained as a deoxidizer, but if it exceeds 0.07 mass%, the cleanliness of the steel is lowered and the HIC resistance is deteriorated, so the Al content is preferably within this range. More preferably, it is 0.01-0.07 mass%.
Cu: 0.5mass%以下とする。
Cuは靭性の改善と強度上昇に有効な元素である。その効果を得るためには、0.1mass%以上含有することが好ましいが、0.5mass%を超えると溶接性が劣化するため、含有する場合はCu含有量をこの範囲に規定する。
Cu: 0.5 mass% or less.
Cu is an element effective in improving toughness and increasing strength. In order to acquire the effect, it is preferable to contain 0.1 mass% or more, but when it exceeds 0.5 mass%, weldability will deteriorate, and when it contains, Cu content is prescribed | regulated to this range.
Ni: 0.5mass%以下とする。
Niは靭性の改善と強度上昇に有効な元素である。その効果を得るためには、0.1mass%以上含有することが好ましいが、0.5mass%を超えると耐HIC特性が劣化するため、含有する場合はNi含有量をこの範囲に規定する。
Ni: 0.5 mass% or less.
Ni is an element effective in improving toughness and increasing strength. In order to acquire the effect, it is preferable to contain 0.1 mass% or more, but when it exceeds 0.5 mass%, the HIC resistance is deteriorated.
Cr: 0.5mass%以下とする。
CrはMnと同様に強度上昇に有効な元素である。その効果を得るためには、0.1mass%以上含有することが好ましいが、0.5mass%を超えると溶接性が劣化しやすくなるため、含有する場合はCr含有量をこの範囲に規定する。
Cr: 0.5 mass% or less.
Cr, like Mn, is an effective element for increasing the strength. In order to acquire the effect, it is preferable to contain 0.1 mass% or more, but if it exceeds 0.5 mass%, the weldability tends to deteriorate, so when it is contained, the Cr content is specified within this range.
Ca: 0.0005〜0.003mass%とする。
既に述べた通り、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、0.0005mass%未満ではその効果が十分ではなく、0.003mass%を超えると鋼の清浄度の低下により耐HIC特性を劣化させることがあるのでこの範囲に規定する。
Ca: 0.0005 to 0.003 mass%.
As already mentioned, Ca is an element effective for improving HIC resistance by controlling the form of sulfide inclusions, but if it is less than 0.0005 mass%, the effect is not sufficient, and if it exceeds 0.003 mass%, the cleanliness of the steel The HIC resistance may be deteriorated due to a decrease in the thickness, so this range is specified.
N: 0.007 mass%以下とする。
Nが0.007mass%を超えるとTi含有量を0.03mass%まで増やしてもTiで固定できない固溶Nが靭性に悪影響を及ぼすため、この範囲に規定する。また、析出強化に有効なTiを含む微細複合炭化物を生成させるためにスラブ加熱段階で固溶Tiを確保するという観点から、N含有量は少ないほど良い。
N: 0.007 mass% or less.
If N exceeds 0.007 mass%, solid solution N that cannot be fixed with Ti even if the Ti content is increased to 0.03 mass% adversely affects toughness. Further, from the viewpoint of securing solid solution Ti in the slab heating stage in order to produce fine composite carbide containing Ti effective for precipitation strengthening, the smaller the N content, the better.
Nb: 0.04〜0.05 mass%とする。
Nbは組織の微細粒化により靭性を向上させるが、同時に、本発明では強度上昇に特に有効な大きさ20nm未満の炭化物形成元素としても機能する。しかし、0.04 mass%未満では、析出強化の効果が不十分となりやすい。また、0.05mass%を超えるとスラブ加熱時に粗大析出して母材靭性の劣化を招く他、HAZに粗大析出してHAZ靭性の劣化を招く。よって、Nb含有量はこの範囲に規定する。なお、後述するTi、Mo、Vと共に複合炭化物を形成することもある。
Nb: 0.04 to 0.05 mass%.
Nb improves toughness by making the structure finer, but at the same time, in the present invention, it functions as a carbide-forming element having a size of less than 20 nm that is particularly effective for increasing the strength. However, if it is less than 0.04 mass%, the effect of precipitation strengthening tends to be insufficient. On the other hand, if it exceeds 0.05 mass%, it coarsely precipitates during slab heating and causes deterioration of the base metal toughness, and also coarsely precipitates in the HAZ and causes deterioration of HAZ toughness. Therefore, Nb content is prescribed | regulated to this range. A composite carbide may be formed together with Ti, Mo, and V described later.
Ti: 0.01〜0.03 mass%とする。
前述したNbに対するTiの析出牽引効果を利用するためには、0.01mass%未満ではその効果が不十分で、0.03mass%を超えるとHAZに粗大析出してHAZ靭性の劣化を招くため、Ti含有量はこの範囲が好ましい。更に好ましくは、0.015〜0.025mass%である。
Ti: 0.01 to 0.03 mass%.
In order to use the Ti precipitation traction effect on Nb as described above, the effect is insufficient if it is less than 0.01 mass%, and if it exceeds 0.03 mass%, it coarsely precipitates in HAZ and causes deterioration of HAZ toughness. The amount is preferably within this range. More preferably, it is 0.015-0.025 mass%.
Mo: 0.05〜0.2 mass%が好ましい。
Moは、TiやNbなどと共に微細な複合炭化物を形成して強度上昇に寄与する。また、Moの炭化物は、TiやNbの炭化物に比べて成長速度が遅いため、複合炭化物が粗大化するのを抑制する作用を有する。0.05mass%以上含有す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、析出強化に有効に働く。しかし、0.5 mass%を超えて含有するとマルテンサイトなどの硬質相を形成して耐HIC特性が劣化しやすくなる。更に、HAZ靭性の観点から、0.2mass%以下が好ましいため、含有する場合にはMo含有量をこの範囲に規定する。
Mo: 0.05 to 0.2 mass% is preferable.
Mo forms a fine composite carbide with Ti, Nb, etc., and contributes to an increase in strength. In addition, Mo carbide has a growth rate slower than that of Ti or Nb carbide, and thus has an action of suppressing the coarsening of the composite carbide. Containing 0.05 mass% or more effectively works for precipitation strengthening while suppressing pearlite transformation during cooling after hot rolling. However, if the content exceeds 0.5 mass%, a hard phase such as martensite is formed and the HIC resistance tends to deteriorate. Furthermore, from the viewpoint of HAZ toughness, 0.2 mass% or less is preferable, and when it is contained, the Mo content is specified within this range.
V: 0.005〜0.1mass%が好ましい。
VはMoと同様にTiやNbと共に微細な複合炭化物を形成して強度上昇に寄与する。また、Vの炭化物は、Moの炭化物と同様に複合炭化物が粗大化するのを抑制する作用を有する。しかし、0.005mass%では効果が得られず、0.1mass%を超えるとHAZ靭性の劣化を招きやすくなるため、含有する場合にはこの範囲に規定する。
V: 0.005 to 0.1 mass% is preferable.
V, like Mo, forms fine composite carbides with Ti and Nb and contributes to an increase in strength. Further, the carbide of V has an action of suppressing the coarsening of the composite carbide, like the carbide of Mo. However, if 0.005 mass%, the effect cannot be obtained, and if it exceeds 0.1 mass%, the HAZ toughness tends to be deteriorated.
尚、Ti、Nb、Mo、Vは、その全量がMC型(金属元素とCのモル比が等量)の析出物を形成できるように、その合計のモル濃度がCのモル濃度以下に納まるように含有する。その際、MoとVは、少なくとも何れか1つが含有されていればよい。 In addition, Ti, Nb, Mo, and V are all in a total molar concentration within the molar concentration of C so as to form precipitates of MC type (equal molar ratio of metal element to C). Contained. At this time, it is sufficient that at least one of Mo and V is contained.
なお、上記以外の残部はFe及び不可避的不純物からなる。不可避的不純物として、例えば、Oは非金属介在物を形成し品質、特に靭性に悪影響を及ぼすため、0.003mass%以下に低減するのが好ましい。また、本発明では、本発明の作用効果を害さない微量元素として、Pを0.01mass%以下、Sを0.002mass%以下の範囲で含有してもよい。 The balance other than the above consists of Fe and inevitable impurities. As an unavoidable impurity, for example, O forms non-metallic inclusions and adversely affects quality, particularly toughness, so it is preferably reduced to 0.003 mass% or less. Moreover, in this invention, you may contain P in 0.01 mass% or less and S in the range of 0.002 mass% or less as a trace element which does not impair the effect of this invention.
また、上記に加えて、TiはN含有量の3.4倍に0.003mass%を加えた値以上含有する。Tiは窒化物を形成しやすい元素のため、N含有量の3.4倍以下の含有ではスラブ加熱段階で、そのほぼ全量が窒化物を形成して、後工程で炭化物を析出させるための固溶Tiを確保できなくなる。そのため、スラブ加熱段階でNの全量がTiを窒化した場合でも、大きさ20nm未満の析出物だけでTiを10mass ppm以上含むように析出させるためには、固溶Tiを少なくともN含有量の3.4倍に30mass ppm(0.003mass%)を加えた値は確保することが好ましいために下限をこのように規定する。 In addition to the above, Ti contains at least 3.4 times the N content plus 0.003 mass%. Ti is an element that tends to form nitrides, so if it contains less than 3.4 times the N content, solid solution Ti will form nitrides in the slab heating stage, and precipitate carbides in the subsequent process. Cannot be secured. Therefore, even when the entire amount of N is nitrided in the slab heating stage, in order to precipitate Ti so as to contain Ti of 10 mass ppm or more with only precipitates having a size of less than 20 nm, solute Ti must be at least N having a N content of 3.4. Since it is preferable to secure a value obtained by adding 30 mass ppm (0.003 mass%) to the double, the lower limit is defined in this way.
2)炭素当量Ceqを0.41未満とする。
既に述べたとおり、炭素当量Ceq は高強度化のための重要因子だが、0.41以上では表層や円周溶接部の最高硬さが高くなって靭性を安定的に確保できなくなるため、0.41未満に規定する。
なお、炭素当量Ceqは、C+Mn/6+Cu/15+Ni/15+Cr/5+Mo/5+V/5(各元素はmass%)にて計算される。
2) The carbon equivalent Ceq is less than 0.41.
As already mentioned, the carbon equivalent Ceq is an important factor for increasing the strength, but if it is 0.41 or higher, the maximum hardness of the surface layer and circumferential welds will be high, and it will not be possible to stably secure toughness. To do.
The carbon equivalent Ceq is calculated by C + Mn / 6 + Cu / 15 + Ni / 15 + Cr / 5 + Mo / 5 + V / 5 (each element is mass%).
3)次に、本発明の高強度鋼材の析出物・組織について説明する。 3) Next, the precipitate and structure of the high-strength steel material of the present invention will be described.
大きさ20nm未満の析出物中のTiは鋼材全体に対し10mass ppm以上含有することとする。析出強化に有効なこのような微細析出物を形成させたとしても、その量がTi含有量で鋼材全体に対し10mass ppm未満では十分な効果が得られにくいため、このように規定する。 Ti in precipitates with a size of less than 20 nm should be contained in an amount of 10 mass ppm or more with respect to the entire steel material. Even if such fine precipitates effective for precipitation strengthening are formed, if the amount is less than 10 mass ppm with respect to the whole steel material with a Ti content, it is difficult to obtain a sufficient effect, so it is specified in this way.
加えて、大きさ20nm未満の析出物中のNbは鋼材全体に対し140mass ppm以上とする。一般に、析出強化量は、析出粒子の大きさとその個数密度で決まり、大きさが小さく、個数密度が高いほど効果的であることが知られている。そのため、鋼材全体に対し10mass ppm程度のTiだけで、一定量以上の析出強化量を実現しようとすると、Ti析出物を極めて微細かつ高密度に形成させなければならないため、その実現が困難な場合がある。利用する析出物を複合析出物の形にして、その形成にNbを利用できれば、同析出物の個数密度を高めて一定量以上の析出強化量を実現するのは格段に容易になる。本発明者の研究によれば、その際に効果的な大きさ20nm未満の析出物中のNb含有量は鋼材全体に対し140mass ppm以上であることが好ましいことが知見されたため、このように規定する。 In addition, Nb in precipitates with a size of less than 20 nm is 140 mass ppm or more with respect to the entire steel material. Generally, the precipitation strengthening amount is determined by the size of the precipitated particles and the number density thereof, and it is known that the smaller the size and the higher the number density, the more effective the precipitation strengthening amount. Therefore, if it is difficult to realize the precipitation strengthening amount of a certain amount or more with only Ti of about 10 mass ppm for the whole steel material, the Ti precipitates must be formed very finely and with high density. There is. If the precipitate to be used is formed into a composite precipitate and Nb can be used to form the precipitate, it is much easier to increase the number density of the precipitate and realize a precipitation strengthening amount of a certain amount or more. According to the inventor's research, it was found that the Nb content in the precipitate having an effective size of less than 20 nm at that time is preferably 140 mass ppm or more with respect to the entire steel material. To do.
更に、大きさ20nm未満の析出物中には鋼材全体に対し50mass ppm以上のMo、および/または、鋼材全体に対し5mass ppm以上のVが含有されていることが好ましい。上記複合析出物の形成に、TiとNbだけでなく、MoやVも利用できれば、その個数密度を高めて一定量以上の析出強化量を実現するのは更に容易になる。また、MoやVの炭化物は、TiやNbの炭化物に比べて成長速度が遅いため、複合析出物が粗大化するのを抑制する作用を有する。そのため、MoやVは、複合析出粒子の個数密度を高めて析出強化量を向上させるのに有効であり、また複合析出物がHAZで一旦溶解する場合にも、それが再析出の際に粗大化して靭性を損なうのを抑制する作用も有する。本発明者の研究によれば、その際に効果的な大きさ20nm未満の析出物中のMoは鋼材全体に対し50mass ppm以上、Vは鋼材全体に対し5 mass ppm以上であることが好ましいため、このように規定する。 Furthermore, it is preferable that the precipitates having a size of less than 20 nm contain 50 mass ppm or more of Mo with respect to the whole steel material and / or 5 mass ppm or more of V with respect to the whole steel material. If not only Ti and Nb but also Mo and V can be used for the formation of the composite precipitate, it becomes easier to increase the number density and realize a precipitation strengthening amount of a certain amount or more. In addition, since carbides of Mo and V have a slower growth rate than Ti and Nb carbides, they have an action of suppressing the coarsening of the composite precipitate. Therefore, Mo and V are effective in increasing the number density of the composite precipitate particles and improving the precipitation strengthening amount. Even when the composite precipitate is once dissolved in HAZ, it is coarse during reprecipitation. It also has an effect of suppressing the deterioration of toughness. According to the inventor's research, Mo in precipitates having an effective size of less than 20 nm is preferably 50 mass ppm or more with respect to the entire steel material, and V is preferably 5 mass ppm or more with respect to the entire steel material. Stipulate in this way.
大きさ20nm未満の析出物に含まれるTi, Nb, Mo, Vの含有量は、以下の方法により確認することができる。 The content of Ti, Nb, Mo, V contained in the precipitate having a size of less than 20 nm can be confirmed by the following method.
試料を電解液中で所定量電解した後、試料片を電解液から取り出して分散性を有する溶液中に浸漬する。次いで、この溶液中に含まれる析出物を、孔径20nmのフィルタを用いてろ過する。この孔径20nmのフィルタをろ液と共に通過した析出物が大きさ20nm未満である。次いで、ろ過後のろ液に対して、誘導結合プラズマ(ICP)発光分光分析法、ICP質量分析法、および原子吸光分析法等から適宜選択して分析し、大きさ20nm未満での析出物におけるTi, Nb, Mo, Vの鋼材全体に対する含有量を求める。 After the sample is electrolyzed in a predetermined amount in the electrolytic solution, the sample piece is taken out of the electrolytic solution and immersed in a solution having dispersibility. Next, the precipitate contained in the solution is filtered using a filter having a pore diameter of 20 nm. The precipitate that has passed through the filter having a pore diameter of 20 nm together with the filtrate has a size of less than 20 nm. Next, the filtrate after filtration is analyzed by appropriately selecting from inductively coupled plasma (ICP) emission spectroscopy, ICP mass spectrometry, atomic absorption spectrometry, etc. Obtain the content of Ti, Nb, Mo, and V for the entire steel.
大きさ20nm未満の析出物のTi含有量を鋼材全体に対し10mass ppm以上にする方法としては、例えば、以下のような方法がある。
まず、上記成分で溶製した鋼を1050〜1250℃の範囲のスラブ加熱温度で加熱して溶体化した後、所定の板厚までAr3 温度以上の温度域(オーステナイト単相の温度域)で圧延し、同温度域からベイナイト変態域まで5℃/sec以上の冷却速度で300〜600℃まで加速冷却する。これによって徐冷した場合に高温域から粗大な析出物が成長することを抑制する。次いで、再加熱設備までの搬送の間、冷却停止温度で0.5〜3分間放冷した後、0.5〜2.0℃/secの昇温速度で最高到達温度600〜700℃まで再加熱し、すぐに(すなわち、保持時間を設けずに)空冷時の冷却速度以上の冷却速度で冷却する。この再加熱によって、金属組織をフェライトとベイナイトの2相組織とし、モル比で比較してもTiよりNbリッチな大きさ20nm未満の複合析出物をこれらの中に分散析出させることができる。このTiの析出牽引効果の本質は充分明らかになっていないが、Tiが炭化物(もしくは炭窒化物)を形成する熱力学的な駆動力がNb、Mo、Vよりも大きく、析出したTi炭化物が複合析出物の析出核として機能するため、Nb、Mo、Vが個々に析出する際に乗り越えなければならないエネルギー障壁(核生成障壁)が不要になって複合析出が促進されるためだと推定される。
なお、オーステナイト単相の温度域で圧延するとは、圧延終了時の温度がAr3温度(フェライト変態開始温度)以上であることを意味し、Ar3=910-310C-80Mn-20Cu-15Cr-55Ni-80Mo から求めることができる。ここでAr3温度の単位は℃で、成分濃度はいずれもmass%である。ベイナイト変態域は300〜600℃である。
再加熱時の昇温速度が0.5℃/sec未満では目的の到達温度に達するまでに長時間を要するため、製造効率が悪く、また、パーライト変態が生じて上記複合析出物の析出と競合するセメンタイトの析出が起こるため、微細複合析出物を必要量確保することができない。また、昇温速度が2.0℃/sec超えでも微細複合析出物を必要量確保することが困難である。これは微細析出物の析出速度がこの昇温速度に追従できないためと推定される。
大きさ20nm未満の微細析出物の粗大化による強度低下を抑制する観点から再加熱温度で保持時間を設定する必要はない。再加熱装置としては、鋼板の急速加熱が可能なガス燃焼炉や誘導加熱炉を用いることが好ましい。
Examples of a method for setting the Ti content of precipitates having a size of less than 20 nm to 10 mass ppm or more with respect to the entire steel material include the following methods.
First, the steel melted with the above components is heated at a slab heating temperature in the range of 1050 to 1250 ° C to form a solution, and then rolled in a temperature range above the Ar3 temperature (austenite single-phase temperature range) up to a predetermined thickness. And accelerated cooling from 300 to 600 ° C. at a cooling rate of 5 ° C./sec or more from the same temperature range to the bainite transformation range. This suppresses the growth of coarse precipitates from the high temperature range when it is gradually cooled. Next, during transportation to the reheating equipment, after cooling for 0.5 to 3 minutes at the cooling stop temperature, reheating to a maximum temperature of 600 to 700 ° C at a heating rate of 0.5 to 2.0 ° C / sec, immediately ( In other words, cooling is performed at a cooling rate equal to or higher than the cooling rate during air cooling (without providing a holding time). By this reheating, the metal structure becomes a two-phase structure of ferrite and bainite, and composite precipitates having a size less than 20 nm, which is Nb richer than Ti, can be dispersed and precipitated even in a molar ratio. Although the essence of this Ti precipitation traction effect has not been fully clarified, the thermodynamic driving force by which Ti forms carbides (or carbonitrides) is larger than that of Nb, Mo, and V. It is presumed that because it functions as a precipitation nucleus for composite precipitates, it eliminates the need for energy barriers (nucleation barriers) that must be overcome when Nb, Mo, and V individually precipitate, thereby promoting composite precipitation. The
In addition, rolling in the temperature range of the austenite single phase means that the temperature at the end of rolling is higher than the Ar3 temperature (ferrite transformation start temperature), Ar3 = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo Can be obtained from Here, the unit of Ar3 temperature is ° C., and the concentration of each component is mass%. The bainite transformation range is 300-600 ° C.
When the heating rate during reheating is less than 0.5 ° C / sec, it takes a long time to reach the target temperature, so that the production efficiency is poor, and pearlite transformation occurs and competes with the precipitation of the composite precipitate. Therefore, the required amount of fine composite precipitates cannot be ensured. In addition, it is difficult to secure the necessary amount of fine composite precipitates even when the rate of temperature rise exceeds 2.0 ° C./sec. This is presumably because the precipitation rate of fine precipitates cannot follow this rate of temperature increase.
It is not necessary to set the holding time at the reheating temperature from the viewpoint of suppressing strength reduction due to coarsening of fine precipitates having a size of less than 20 nm. As the reheating device, it is preferable to use a gas combustion furnace or an induction heating furnace capable of rapid heating of the steel sheet.
つぎに、本発明の実施例について以下に説明する。 Next, examples of the present invention will be described below.
表1に示す成分の試験材(符号A〜K)を溶解してスラブとし、これを図1に示す圧延条件ならびに熱処理条件で処理した。まず、スラブ加熱を1150℃で2時間おこなった。その後、スラブを熱間圧延機により950℃で75%圧延し、更に860℃で仕上げて板厚を約20mmとした後、790℃から水冷型の加速冷却設備を用いて冷却速度30℃/secで500℃まで加速冷却した。しかる後、試験材の一部はそのまま空冷した(熱処理条件1:比較例)。また、残りの試験片は、析出強化をおこなうため、500℃で1分間放置した後、ガス燃焼炉を用いて昇温速度1℃/secで650℃まで再加熱し、しかる後、空冷した(熱処理条件2:本発明例)。 The test materials (signs A to K) having the components shown in Table 1 were dissolved to form slabs, which were processed under the rolling conditions and heat treatment conditions shown in FIG. First, slab heating was performed at 1150 ° C. for 2 hours. After that, the slab was rolled 75% at 950 ° C with a hot rolling mill and finished at 860 ° C to a plate thickness of about 20mm. Then, the cooling rate was 30 ° C / sec using a water-cooled accelerated cooling facility from 790 ° C. And accelerated cooling to 500 ° C. Thereafter, a part of the test material was air-cooled as it was (heat treatment condition 1: comparative example). The remaining test pieces were left to stand at 500 ° C. for 1 minute for precipitation strengthening, and then reheated to 650 ° C. at a heating rate of 1 ° C./sec using a gas combustion furnace, and then air-cooled ( Heat treatment condition 2: Example of the present invention).
得られた鋼材に対して、以下に示す方法で各試験材に含まれる大きさ20nm未満の析出物に含まれる各元素含有量(表1の成分にFeを加えた合計を100mass%とした場合の、Ti、Nb、Mo、Vの含有量)、引張強度、シャルピー衝撃試験での破面遷移温度(vTrs)、耐HIC性を求め評価した。 With respect to the obtained steel materials, the content of each element contained in precipitates of less than 20 nm in size contained in each test material by the method shown below (when the total of the components in Table 1 plus Fe is 100 mass% Of Ti, Nb, Mo, and V), tensile strength, fracture surface transition temperature in Charpy impact test (vTrs), and HIC resistance were evaluated.
イ)大きさ20nm未満の析出物に含まれるTi、Nb、MoおよびVの鋼材全体に対する含有量の測定
上記により得られた鋼材を適当な大きさに切断し、10%AA系電解液(10vol%アセチルアセトン-1mass%塩化テトラメチルアンモニウム-メタノール)中で、約0.2gを電流密度20mA/cm2で定電流電解した。
電解後の、表面に析出物が付着している試料片を電解液から取り出して、ヘキサメタリン酸ナトリウム水溶液(500mg/l)(以下、SHMP水溶液と称す)中に浸漬し、超音波振動を付与して、析出物を試料片から剥離しSHMP水溶液中に抽出した。次いで、析出物を含むSHMP水溶液を、孔径20nmのフィルタを用いてろ過し、ろ過後のろ液に対してICP発光分光分析装置を用いて分析し、ろ液中のTiとNbとMoとVの絶対量を測定した。次いで、TiとNbとMoとVの絶対量を電解重量で除して、大きさ20nm未満の析出物に含まれるTi、Nb、MoおよびVの含有量を得た。なお、電解重量は、析出物剥離後の試料に対して重量を測定し、電解前の試料重量から差し引くことで求めた。
B) Measurement of the content of Ti, Nb, Mo, and V contained in precipitates with a size of less than 20 nm for the entire steel material The above-obtained steel material was cut to an appropriate size, and 10% AA electrolyte (10 vol. % Acetylacetone-1 mass% tetramethylammonium chloride-methanol) was subjected to constant current electrolysis at a current density of 20 mA / cm 2 .
After the electrolysis, remove the sample piece with deposits on the surface from the electrolyte and immerse it in an aqueous solution of sodium hexametaphosphate (500 mg / l) (hereinafter referred to as the SHMP aqueous solution) to apply ultrasonic vibration. The precipitate was peeled off from the sample piece and extracted into an aqueous SHMP solution. Next, the SHMP aqueous solution containing the precipitate is filtered using a filter with a pore size of 20 nm, and the filtrate after filtration is analyzed using an ICP emission spectroscopic analyzer. Ti, Nb, Mo, and V in the filtrate are analyzed. The absolute amount of was measured. Next, the absolute amounts of Ti, Nb, Mo, and V were divided by the electrolytic weight to obtain the contents of Ti, Nb, Mo, and V contained in the precipitate having a size of less than 20 nm. In addition, the electrolysis weight was calculated | required by measuring a weight with respect to the sample after deposit peeling, and subtracting from the sample weight before electrolysis.
ロ)引張強度
圧延垂直方向の全厚試験片を引張試験片として測定した。熱処理2の引張強度が熱処理1の値を30MPa以上上回る場合に析出強化が効果的に実現されていると判断した。なお、ΔTSは、同じ鋼種(符号)における熱処理2と熱処理1との引張強度の差である。
また、この析出強化に有効に働いた大きさ20nm未満の析出物がMC型であることは、抽出した析出物の電子回折パターンによって確認した。また、その金属成分中のNbのモル比がTiよりもリッチであることは、上記の含有率測定法によって確認した。一例として、上記の方法で測定した符号Bの大きさ20nm未満の析出物の平均組成を示すと、(Ti0.07Nb0.43Mo0.28V0.22)Cである。
B) Tensile strength A full thickness test piece in the vertical direction of rolling was measured as a tensile test piece. It was judged that precipitation strengthening was effectively realized when the tensile strength of
Further, it was confirmed by the electron diffraction pattern of the extracted precipitate that the precipitate having a size of less than 20 nm that worked effectively for the precipitation strengthening was MC type. In addition, it was confirmed by the content ratio measurement method that the molar ratio of Nb in the metal component was richer than Ti. As an example, the average composition of precipitates having a size of code B less than 20 nm measured by the above method is (Ti 0.07 Nb 0.43 Mo 0.28 V 0.22 ) C.
ハ)vTrsの測定
図2に示すHAZを再現した熱処理を実施した上で、シャルピー衝撃試験で破面遷移温度(vTrs)を測定した。この温度が10℃未満の場合に合格と判断した。
C) Measurement of vTrs After heat treatment reproducing HAZ shown in Fig. 2, fracture surface transition temperature (vTrs) was measured by Charpy impact test. When this temperature was less than 10 ° C, it was judged to be acceptable.
ニ)耐HIC特性の測定
耐HIC特性はNACE Standard TM-02-84に準じたpH3.0, 0.1atm H2Sのライトサワー環境で浸漬時間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○とした。
D) Measurement of HIC resistance HIC resistance was measured in a light sour environment of pH 3.0, 0.1 atm H 2 S according to NACE Standard TM-02-84. The case was judged to be good because it was judged to have good HIC resistance.
以上により得られた結果を表2に示す。 The results obtained as described above are shown in Table 2.
表2より、本発明例では、耐HIC特性および溶接熱影響部の靭性を劣化させることなく比較例と比べて強度が30MPa以上上回り、API規格X80グレード以上の強度を確保している。すなわち、耐HIC特性が劣化しない低合金成分で、API規格X65グレード以上の強度と溶接熱影響部の靭性を両立させていることがわかる。 From Table 2, the strength of the present invention example is 30 MPa or more higher than that of the comparative example and the strength of API standard X80 grade or higher is ensured without degrading the HIC resistance and the toughness of the heat affected zone. That is, it can be seen that the low alloy component that does not deteriorate the HIC resistance property satisfies both the strength of API standard X65 grade or higher and the toughness of the weld heat affected zone.
図3は、表1の引張強度を炭素当量Ceqに対してプロットしたものである。20nm未満の析出物で析出強化された本発明例では炭素当量0.41未満のCeqで引張強度が効果的に改善されていることが明らかである。本発明が対象とするAPI X80の引張強度TSが650MPa以上であることから、熱処理条件2の方法が有効であることがわかる。すなわち、大きさ20nm未満の析出物のTi含有量が鋼材全体に対して10mass ppm以上でNb含有量が鋼材全体に対して140mass ppm以上であるものは引張強度が高くなっている。なお、熱処理条件2で得られたものは、HAZ靭性も良好である。
FIG. 3 is a plot of the tensile strength in Table 1 against the carbon equivalent Ceq. It is clear that the tensile strength is effectively improved with the Ceq of less than 0.41 carbon equivalent in the inventive examples where precipitation strengthening is performed with precipitates of less than 20 nm. Since the tensile strength TS of API X80 targeted by the present invention is 650 MPa or more, it can be seen that the method of
図4は、表2のvTrs(再現HAZ)をTi含有量に対してプロットした結果である。Ti含有量が0.03mass%までは、vTrsを良好な範囲に抑えられることがわかる。 FIG. 4 shows the results of plotting vTrs (reproduced HAZ) in Table 2 against Ti content. It can be seen that when the Ti content is 0.03 mass%, vTrs can be suppressed within a favorable range.
本発明の鋼板は、耐HIC特性が必要とされる環境下に使用され、API規格X65グレード以上の強度および溶接熱影響部の靭性を両立させるラインパイプ用鋼材として好適に使用できる。 The steel sheet of the present invention is used in an environment where HIC resistance is required, and can be suitably used as a steel material for a line pipe that achieves both the strength of API standard X65 grade or higher and the toughness of the heat affected zone.
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