JP6149368B2 - Manufacturing method of high-tensile steel plate with excellent delayed fracture resistance - Google Patents

Manufacturing method of high-tensile steel plate with excellent delayed fracture resistance Download PDF

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JP6149368B2
JP6149368B2 JP2012213638A JP2012213638A JP6149368B2 JP 6149368 B2 JP6149368 B2 JP 6149368B2 JP 2012213638 A JP2012213638 A JP 2012213638A JP 2012213638 A JP2012213638 A JP 2012213638A JP 6149368 B2 JP6149368 B2 JP 6149368B2
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JP2014029003A (en
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茂樹 木津谷
茂樹 木津谷
浩文 大坪
浩文 大坪
謙次 林
謙次 林
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JFE Steel Corp
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本発明は、耐遅れ破壊特性に優れた高張力鋼板の製造方法に関し、特に引張強さが780MPa以上の高張力鋼材として好適なものに関する。   The present invention relates to a method for producing a high-strength steel sheet having excellent delayed fracture resistance, and particularly relates to a suitable high-tensile steel material having a tensile strength of 780 MPa or more.

近年、建設産業機械、タンク、ペンストック、ラインパイプ等の鋼材使用分野では、構造物の大型化、軽量化を背景として、使用する鋼材の高強度化が指向されると共に鋼材使用量が急激に増加している。   In recent years, in the field of using steel materials such as construction industrial machines, tanks, penstocks, line pipes, etc., against the backdrop of increasing the size and weight of structures, the strength of steel materials to be used has been increasing and the amount of steel materials used has rapidly increased. It has increased.

このような需要の増大に対しては、特許文献1、2、3には、焼戻し処理を省略した780MPa級高張力厚鋼板の製造方法が提案されている。   In response to such an increase in demand, Patent Documents 1, 2, and 3 propose a method of manufacturing a 780 MPa class high-tensile thick steel plate that omits the tempering treatment.

また、さらに高強度鋼の需要に対応するため、特許文献4には、焼戻し処理を省略した降伏強度885MPa以上の高張力鋼板の製造方法が提案されている。   In order to further meet the demand for high-strength steel, Patent Document 4 proposes a method for manufacturing a high-strength steel sheet having a yield strength of 885 MPa or more, in which tempering treatment is omitted.

特開2007−277622号公報JP 2007-277622 A 特開2007−277623号公報JP 2007-277623 A 特開2009−263772号公報JP 2009-263774 A 特開2011−012315号公報JP 2011-012315 A

特許文献1〜3には「焼戻し処理を省略した引張強度780MPa級高強度鋼板の製造方法として、未再結晶域圧延を完了後、700℃以上から冷却速度が8〜80℃/sとなる加速冷却を開始し、室温〜350℃で加速冷却を停止することが記載されており、この製造方法は、自己焼鈍による不均一な炭化物の生成が起こりやすく、耐遅れ破壊特性の観点からは問題がある。また、焼入れままの強度を調整するためにCを低減する必要があり、必要な特性を満足するためには大量の合金元素を添加する必要があり、経済性が低下するという問題もある。   In Patent Documents 1 to 3, “as a method for producing a high strength steel sheet having a tensile strength of 780 MPa class without tempering treatment, an acceleration at which the cooling rate is 8 to 80 ° C./s from 700 ° C. or higher after completion of non-recrystallization zone rolling. It is described that cooling is started and accelerated cooling is stopped at room temperature to 350 ° C., and this production method is liable to generate non-uniform carbides by self-annealing, and there is a problem from the viewpoint of delayed fracture resistance. In addition, it is necessary to reduce C in order to adjust the as-quenched strength, and in order to satisfy the necessary characteristics, it is necessary to add a large amount of alloying elements, and there is a problem that economical efficiency is lowered. .

特許文献4には「焼戻し処理を省略した降伏強度885MPa以上の高強度鋼板の製造方法として、未再結晶域圧延を完了後Ar以上の温度から平均冷却速度10℃/s以上で加速冷却を行い200〜400℃で冷却を停止し、その後20℃/min以下で冷却する」ことが記載されており、この製造方法は、特許文献1〜3と同様に、自己焼鈍による不均一な炭化物の生成が起こりやすく、耐遅れ破壊特性の観点からは問題がある。
また、オートエージングにより強度を確保しているため、低温靭性も十分とは言い難い。
In Patent Document 4, “as a method for producing a high-strength steel sheet having a yield strength of 885 MPa or more without tempering treatment, accelerated cooling is performed at an average cooling rate of 10 ° C./s or more from a temperature of Ar 3 or more after completing non-recrystallization zone rolling. The cooling is stopped at 200 to 400 ° C. and then cooled at 20 ° C./min or less ”, and this manufacturing method is similar to Patent Documents 1 to 3 in that non-uniform carbides by self-annealing are described. It tends to occur and is problematic in terms of delayed fracture resistance.
Moreover, since the strength is secured by auto-aging, it is difficult to say that the low temperature toughness is sufficient.

本発明はかかる事情に鑑みてなされたものであり、引張強度が780MPa以上で、従来の鋼材より耐遅れ破壊特性に優れた高張力鋼板の製造方法を提供することを目的とする。   The present invention has been made in view of such circumstances, and an object of the present invention is to provide a method for producing a high-tensile steel sheet having a tensile strength of 780 MPa or more and superior in delayed fracture resistance to conventional steel materials.

本発明者らは、上記課題を解決するために、780MPa以上の強度と耐遅れ破壊特性を有する鋼板について鋭意研究を重ねた結果、焼戻し処理時における鋼板の板厚方向中心部の昇温速度、焼戻し温度および保持時間を規定することによって、焼戻しによる炭化物の生成、粗大化が効果的に抑制され、拡散性水素の炭化物への集積を抑えることができ、焼戻し加熱時の脱水素により鋼中の拡散性水素量を低減することにより、切断時および使用中の遅れ破壊を抑制することが可能となり、従来鋼よりも耐遅れ破壊特性に優れた高張力鋼板が得られることを知見した。   In order to solve the above-mentioned problems, the present inventors have conducted extensive research on a steel sheet having a strength of 780 MPa or more and delayed fracture resistance, and as a result, the rate of temperature rise at the center of the steel sheet in the thickness direction during tempering, By defining the tempering temperature and holding time, the formation and coarsening of carbides due to tempering can be effectively suppressed, and accumulation of diffusible hydrogen into carbides can be suppressed. It has been found that by reducing the amount of diffusible hydrogen, delayed fracture during cutting and during use can be suppressed, and a high-tensile steel sheet having better delayed fracture resistance than conventional steel can be obtained.

本発明は上記の得られた知見に基づき、更に検討を加えてなされたもので、その要旨は、以下の通りである。   The present invention has been made on the basis of the above-described findings and has been further studied. The summary of the present invention is as follows.

[1]質量%で、C:0.02〜0.25%、Si:0.01〜0.8%、Mn:0.5〜2%、P:0.010%以下、S:0.003%以下、Al:0.005〜0.1%、N:0.0005〜0.008%を含有し、残部がFeおよび不可避的不純物からなる鋼をAc変態点以上に加熱し、未再結晶温度域での累積圧下率を80%以下とする熱間圧延を行い、Ar変態点以上で熱間圧延を終了し、引き続きAr変態点以上から10℃/s以上の冷却速度で250℃以下の温度まで冷却後、1℃/s以上の平均昇温速度で再加熱し、最高到達温度を100〜400℃の範囲とする焼戻し処理を行うことを特徴とする耐遅れ破壊特性に優れた高張力鋼板の製造方法。 [1] By mass%, C: 0.02 to 0.25%, Si: 0.01 to 0.8%, Mn: 0.5 to 2%, P: 0.010% or less, S: 0.0. A steel containing 0.003% or less, Al: 0.005 to 0.1%, N: 0.0005 to 0.008%, with the balance being Fe and unavoidable impurities, is heated to the Ac 3 transformation point or higher, subjected to hot rolling to a cumulative reduction ratio is 80% or less at the recrystallization temperature region, Ar 3 hot rolling finished by transformation point or higher, subsequently Ar 3 transformation point at 10 ° C. / s or more cooling rate from above After cooling to a temperature of 250 ° C. or lower, reheat at an average temperature increase rate of 1 ° C./s or higher, and perform a tempering treatment with a maximum temperature of 100 to 400 ° C. An excellent method for producing high-tensile steel sheets.

[2]前記鋼に、更に、質量%で、Mo:0.01〜1%、Nb:0.001〜0.1%、V:0.001〜0.5%、Ti:0.001〜0.1%、Cu:2%以下、Ni:4%以下、Cr:2%以下、W:2%以下の中から選ばれる一種以上を含有することを特徴とする前記[1]に記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   [2] In addition to the steel, by mass%, Mo: 0.01-1%, Nb: 0.001-0.1%, V: 0.001-0.5%, Ti: 0.001- It contains at least one selected from 0.1%, Cu: 2% or less, Ni: 4% or less, Cr: 2% or less, W: 2% or less. A method for manufacturing high-tensile steel sheets with excellent delayed fracture resistance.

[3]前記鋼に、更に、質量%で、B:0.0003〜0.003%、Ca:0.01%以下、REM:0.02%以下の中から選ばれる一種以上を含有することを特徴とする前記[1]または[2]に記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   [3] The steel further contains at least one kind selected from B: 0.0003 to 0.003%, Ca: 0.01% or less, and REM: 0.02% or less in mass%. The method for producing a high-tensile steel sheet having excellent delayed fracture resistance according to [1] or [2] above.

[4]圧延後の再加熱を圧延機および冷却装置と同一の製造ライン上に設置された加熱装置を用いて行い、焼戻し処理において最高到達温度における保持時間を60s以内とすることを特徴とする前記[1]乃至[3]の何れかに記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   [4] Reheating after rolling is performed using a heating device installed on the same production line as the rolling mill and the cooling device, and the holding time at the highest temperature in the tempering process is set to 60 s or less. The method for producing a high-tensile steel sheet having excellent delayed fracture resistance according to any one of [1] to [3].

[5]前記鋼が、質量%で、Nb:0.001〜0.1%を含有し、未再結晶温度域での累積圧下率を40%以上、熱間圧延終了温度を860℃以下とすることを特徴とする請求項1乃至4の何れかに記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   [5] The steel contains, by mass%, Nb: 0.001 to 0.1%, a cumulative rolling reduction in the non-recrystallization temperature range of 40% or more, and a hot rolling end temperature of 860 ° C. or less. The method for producing a high-tensile steel plate having excellent delayed fracture resistance according to any one of claims 1 to 4.

[6]前記高張力鋼板の金属組織はマルテンサイトおよび下部ベイナイトを主体とし、前記金属組織中のラス界面における炭化物の円相当径は100nm以下、炭化物被覆率は30%以下であることを特徴とする前記[1]乃至[5]の何れかに記載の溶接性および耐遅れ破壊特性に優れた高張力鋼板の製造方法。   [6] The metal structure of the high-strength steel sheet is mainly composed of martensite and lower bainite, the equivalent circle diameter of carbide at the lath interface in the metal structure is 100 nm or less, and the carbide coverage is 30% or less. The method for producing a high-tensile steel sheet excellent in weldability and delayed fracture resistance according to any one of [1] to [5].

本発明により、引張強度が780MPa以上の高強度を有するとともに、耐遅れ破壊特性に優れる高張力鋼を安価に安定して製造することができる。   According to the present invention, a high-tensile steel having a high tensile strength of 780 MPa or more and excellent in delayed fracture resistance can be stably produced at a low cost.

以下に本発明の各構成要件の限定理由について説明する。   The reasons for limiting the respective constituent requirements of the present invention will be described below.

1.成分組成について
はじめに、本発明の鋼の成分組成を規定した理由を説明する。なお、成分%は、すべて質量%を意味する。
1. About component composition First, the reason which prescribed | regulated the component composition of the steel of this invention is demonstrated. In addition, all component% means the mass%.

C:0.02〜0.25%
Cは、構造用鋼に求められる強度を得るために必要不可欠な元素であるが、0.02%未満の添加では、十分な強度が得られず、合金元素の大量添加が必要になり溶接性が低下する。0.25%を超えて添加すると、溶接熱影響部のマルテンサイトの生成量が多くなり靭性が低下するため、C量は0.02〜0.25%の範囲とする。好ましくは0.12〜0.22%の範囲である。より好ましくは0.12〜0.18%の範囲である。
C: 0.02-0.25%
C is an indispensable element for obtaining the strength required for structural steel, but if it is added less than 0.02%, sufficient strength cannot be obtained, and a large amount of alloying element is required, so that weldability is achieved. Decreases. If added over 0.25%, the amount of martensite produced in the weld heat affected zone increases and the toughness decreases, so the C content is in the range of 0.02 to 0.25%. Preferably it is 0.12 to 0.22% of range. More preferably, it is 0.12 to 0.18% of range.

Si:0.01〜0.8%
Siは脱酸のために添加するが、0.01%未満の添加では脱酸効果が十分でなく、0.8%を超えて添加すると母材および溶接熱影響部の靭性が顕著に低下するとともに溶接性が著しく低下するため、Si量は0.01%〜0.8%の範囲とする。好ましくは0.1〜0.5%の範囲である。より好ましくは0.2〜0.5%の範囲である。
Si: 0.01 to 0.8%
Si is added for deoxidation, but if less than 0.01% is added, the deoxidation effect is not sufficient, and if added over 0.8%, the toughness of the base metal and the weld heat-affected zone significantly decreases. At the same time, the weldability is remarkably lowered, so the Si content is in the range of 0.01% to 0.8%. Preferably it is 0.1 to 0.5% of range. More preferably, it is 0.2 to 0.5% of range.

Mn:0.5〜2%
Mnは母材強度を確保する観点から添加するが、0.5%未満の添加ではその効果が十分でなく、2%を超えて添加すると、過剰に焼入性を高め、溶接熱影響部の靭性を著しく低下させることから、Mn量は0.5〜2%の範囲とする。好ましくは0.6〜1.6%の範囲である。より好ましくは0.6〜1.3%の範囲である。
Mn: 0.5-2%
Mn is added from the viewpoint of securing the strength of the base metal. However, if it is added in an amount of less than 0.5%, the effect is not sufficient. Since the toughness is remarkably lowered, the amount of Mn is set in the range of 0.5 to 2%. Preferably it is 0.6 to 1.6% of range. More preferably, it is 0.6 to 1.3% of range.

P:0.010%以下
Pは、0.010%を超えて含有すると、母材および溶接熱影響部の靭性を著しく低下させるため、P量は0.010%以下とする。
P: 0.010% or less When P exceeds 0.010%, the toughness of the base metal and the weld heat-affected zone is remarkably reduced, so the P content is 0.010% or less.

S:0.003%以下
Sは、0.003%を超えて含有すると、母材および溶接熱影響部の靭性を顕著に低下させるため、S量は0.003%以下とする。
S: 0.003% or less If S is contained in excess of 0.003%, the toughness of the base metal and the weld heat-affected zone is remarkably reduced. Therefore, the amount of S is made 0.003% or less.

Al:0.005〜0.1%
Alは溶鋼を十分に脱酸するために添加されるが、0.005%未満の添加では脱酸効果が十分でなく、0.1%を超えて添加すると母材中に固溶するAl量が多くなり、母材靭性を低下させるので、Al量は0.005〜0.1%の範囲とする。好ましくは、0.01〜0.06%の範囲である。より好ましくは、0.01〜0.04%の範囲である。
Al: 0.005 to 0.1%
Al is added in order to sufficiently deoxidize molten steel, but if it is added less than 0.005%, the deoxidation effect is not sufficient, and if added over 0.1%, the amount of Al that dissolves in the base metal And the base metal toughness is reduced, so the Al content is in the range of 0.005 to 0.1%. Preferably, it is 0.01 to 0.06% of range. More preferably, it is 0.01 to 0.04% of range.

N:0.0005%〜0.008%
Nは、Tiなどと窒化物を形成することによって組織を微細化し、母材および溶接熱影響部の靭性を向上させる効果を有するために添加する。しかし、0.0005%未満の添加では組織微細化の効果が十分ではなく、一方、0.008%を超えて添加すると、母材中に固溶するN量が増大し、母材靭性が著しく低下し、さらに溶接熱影響部においても粗大な炭窒化物を形成し靭性を低下させるので、N量は0.0005%〜0.008%の範囲とする。好ましくは、0.0010〜0.0060%の範囲である。より好ましくは、0.0010〜0.0040%の範囲である。
N: 0.0005% to 0.008%
N is added in order to refine the structure by forming a nitride with Ti or the like and to improve the toughness of the base material and the weld heat affected zone. However, if the addition is less than 0.0005%, the effect of refining the structure is not sufficient. On the other hand, if the addition exceeds 0.008%, the amount of N dissolved in the base material increases and the base material toughness is remarkably increased. Further, a coarse carbonitride is formed also in the weld heat affected zone and the toughness is lowered. Therefore, the N content is set in the range of 0.0005% to 0.008%. Preferably, it is 0.0010 to 0.0060% of range. More preferably, it is 0.0010 to 0.0040% of range.

以上が本発明の基本化学成分であり、残部はFe及び不可避的不純物からなるが、さらに強度を高める目的でMo、Nb、V、Ti、Cu、Ni、Cr、Wの中から選ばれる1種以上を選択元素として添加してもよい。   The above is the basic chemical component of the present invention, and the balance consists of Fe and unavoidable impurities, but one kind selected from Mo, Nb, V, Ti, Cu, Ni, Cr, and W for the purpose of further increasing the strength. The above may be added as selective elements.

Mo:0.01〜1%
Moは、母材の高強度化に有効な元素であるが、0.01%未満ではその効果が不十分であり、1%を超えて添加すると合金炭化物の析出による硬度の上昇を引き起こし、靭性を低下させるので、Moを添加する場合は、Mo量は0.01〜1%の範囲とするのが好ましい。より好ましくは、0.05〜0.8%の範囲である。
Mo: 0.01 to 1%
Mo is an element effective for increasing the strength of the base material, but its effect is insufficient if it is less than 0.01%, and if added over 1%, it causes an increase in hardness due to precipitation of alloy carbides, and toughness When Mo is added, the amount of Mo is preferably in the range of 0.01 to 1%. More preferably, it is 0.05 to 0.8% of range.

Nb:0.001〜0.1%
Nbは鋼の強化に有効な元素であるとともに適切な圧延条件下で焼戻し後に析出している炭化物を微細化する効果があるが、0.001%未満ではその効果が不十分であり、0.1%を超える添加は母材の靭性を著しく低下させるので、Nbを添加する場合は、Nb量は0.001〜0.1%の範囲とするのが好ましい。より好ましくは、0.001〜0.05%の範囲である。
Nb: 0.001 to 0.1%
Nb is an element effective for strengthening steel and has an effect of refining carbides precipitated after tempering under appropriate rolling conditions. However, if it is less than 0.001%, the effect is insufficient. Since addition exceeding 1% remarkably lowers the toughness of the base material, when Nb is added, the Nb content is preferably in the range of 0.001 to 0.1%. More preferably, it is 0.001 to 0.05% of range.

V:0.001〜0.5%
Vは母材の強度・靭性の向上に効果があり、また、VNとして析出することで固溶Nの低下に有効であるが、0.001%未満ではその効果が不十分であり、0.5%を超えて添加すると硬質なVCの析出により靭性が低下するので、Vを添加する場合は、V量は0.001〜0.5%の範囲とするのが好ましい。より好ましくは、0.01〜0.1%の範囲である。
V: 0.001 to 0.5%
V is effective in improving the strength and toughness of the base metal, and is effective in lowering the solid solution N by being precipitated as VN. However, if it is less than 0.001%, the effect is insufficient. If added over 5%, the toughness decreases due to precipitation of hard VC. Therefore, when V is added, the V content is preferably in the range of 0.001 to 0.5%. More preferably, it is 0.01 to 0.1% of range.

Ti:0.001〜0.1%
Tiは圧延加熱時あるいは溶接時にTiNを生成し、オーステナイトの粗大化を効果的に抑制し、母材および溶接熱影響部の靭性を向上させる。しかし、0.1%を超えて添加すると、Ti窒化物が粗大化し母材および溶接熱影響部の靭性を低下させるので、Tiを添加する場合は、Ti量は0.001〜0.1%の範囲とするのが好ましい。より好ましくは、0.005〜0.020%の範囲である。
Ti: 0.001 to 0.1%
Ti generates TiN at the time of rolling heating or welding, effectively suppressing the coarsening of austenite, and improves the toughness of the base material and the weld heat affected zone. However, if added over 0.1%, the Ti nitride is coarsened and the toughness of the base metal and the weld heat affected zone is reduced. Therefore, when adding Ti, the Ti content is 0.001 to 0.1%. It is preferable to be in the range. More preferably, it is 0.005 to 0.020% of range.

Cu:2%以下
Cuは、低温靭性を損なうことなく鋼の強度の向上が図れるが、2%を超えて添加すると、熱間圧延時に鋼板表面に割れを生じるので、Cuを添加する場合は、Cu量は2%以下とすることが好ましい。より好ましくは、1%以下である。
Cu: 2% or less Cu can improve the strength of the steel without impairing the low-temperature toughness, but if added over 2%, it causes cracking on the surface of the steel sheet during hot rolling, so when adding Cu, The amount of Cu is preferably 2% or less. More preferably, it is 1% or less.

Ni:4%以下
Niは、鋼の強度および溶接熱影響部の靭性を向上させる有益な元素であるが、4%を超えて添加すると、効果が飽和し経済性が劣るため、Niを添加する場合は、Ni量は4%以下とすることが好ましい。より好ましくは、2%以下である。
Ni: 4% or less Ni is a beneficial element that improves the strength of the steel and the toughness of the heat affected zone of the steel, but if added over 4%, the effect is saturated and the economy is inferior, so Ni is added. In this case, the amount of Ni is preferably 4% or less. More preferably, it is 2% or less.

Cr:2%以下
Crは、強度および靭性の向上に有効な元素であるが、2%を超えて添加すると、溶接性が低下するので、Crを添加する場合は、Cr量は2%以下とすることが好ましい。より好ましくは、0.1〜1%の範囲である。
Cr: 2% or less Cr is an element effective for improving the strength and toughness, but if added over 2%, the weldability deteriorates. Therefore, when Cr is added, the Cr amount is 2% or less. It is preferable to do. More preferably, it is 0.1 to 1% of range.

W:2%以下
Wは強度を向上する作用を有している元素であるが、2%を超えて添加すると、溶接性が低下するので、Wを添加する場合は、W量は2%以下とすることが好ましい。より好ましくは、0.05〜2%の範囲である。
W: 2% or less W is an element that has the effect of improving the strength, but if added over 2%, the weldability deteriorates, so when adding W, the amount of W is 2% or less. It is preferable that More preferably, it is 0.05 to 2% of range.

本発明の高張力鋼は、上記組成に加えて、さらに材質を改善する目的でB、Ca、REMの中から選ばれる1種以上を選択元素として添加してもよい。   In addition to the above composition, the high-tensile steel of the present invention may be added with one or more selected from B, Ca, and REM as selective elements for the purpose of further improving the material.

B:0.0003〜0.003%
Bは、オーステナイト粒界に偏析することで粒界からのフェライト変態を抑制し、焼入性を高める効果を有するが、この効果を十分に発揮させるためには0.0003%以上添加することが好ましいが、0.003%を超えて添加すると、炭窒化物として析出し焼入性を低下させ、靭性が低下するので、Bを添加する場合は、B量は0.0003〜0.003%の範囲とするのが好ましい。より好ましくは0.0005〜0.002%の範囲である。
B: 0.0003 to 0.003%
B has the effect of suppressing the ferrite transformation from the grain boundary by segregating at the austenite grain boundary and improving the hardenability, but 0.0003% or more may be added in order to fully exhibit this effect. Preferably, if added over 0.003%, it precipitates as carbonitride, lowering the hardenability and lowering the toughness. Therefore, when adding B, the amount of B is 0.0003 to 0.003%. It is preferable to be in the range. More preferably, it is 0.0005 to 0.002% of range.

Ca:0.01%以下
Caは硫化物系介在物の形態制御に有用な元素である。しかし0.01%を超えて添加すると、清浄度の低下を招くので、Caを添加する場合は、Ca量は0.01%以下とするのが好ましい。より好ましくは0.0005〜0.0020%の範囲である。
Ca: 0.01% or less Ca is an element useful for controlling the form of sulfide inclusions. However, if added over 0.01%, the cleanliness is lowered, so when adding Ca, the Ca content is preferably 0.01% or less. More preferably, it is 0.0005 to 0.0020% of range.

REM:0.02%以下
REMもCaと同様に鋼中で酸化物および硫化物を形成して材質を改善する効果があるが、0.02%を超えて添加しても、その効果が飽和するため、REMを添加する場合は、REM量は0.02%以下とするのが好ましい。より好ましくは0.0005〜0.0020%の範囲である。
REM: 0.02% or less REM also has the effect of improving the quality of the material by forming oxides and sulfides in the steel, similar to Ca, but the effect is saturated even if added over 0.02%. Therefore, when REM is added, the amount of REM is preferably 0.02% or less. More preferably, it is 0.0005 to 0.0020% of range.

2.金属組織について
本発明における金属組織の限定理由について説明する。
2. About Metal Structure The reason for limiting the metal structure in the present invention will be described.

引張強さ780MPa以上の高強度化を図るために金属組織は、マルテンサイトおよび下部ベイナイトを主体とする組織とする必要があり、強度と靭性を両立するためにマルテンサイト+下部ベイナイト組織分率を95%以上とすることが好ましい。5%以下のフェライトや上部ベイナイト、残留γなどは許容する。   In order to increase the tensile strength of 780 MPa or more, the metal structure needs to be a structure mainly composed of martensite and lower bainite. In order to achieve both strength and toughness, the martensite + lower bainite structure fraction should be It is preferably 95% or more. Less than 5% ferrite, upper bainite, residual γ, etc. are allowed.

組織分率の定量化は、ナイタール等の適切な腐食液でエッチングした後に、光学顕微鏡で組織写真を5視野以上撮影し、画像解析によりマルテンサイトと下部ベイナイトの面積率を測定することにより求める。   The structure fraction is quantified by etching with an appropriate corrosive solution such as nital, taking five or more views of the structure photograph with an optical microscope, and measuring the area ratio of martensite and lower bainite by image analysis.

鋼中の拡散性水素量を極力低減して溶接性、耐遅れ破壊特性を向上させる観点から焼もどし処理を施す。   Tempering treatment is performed from the viewpoint of improving weldability and delayed fracture resistance by reducing the amount of diffusible hydrogen in the steel as much as possible.

また、耐遅れ破壊特性の向上には、炭化物の微細分散化が極めて重要である。最も好ましくはマルテンサイトおよび下部ベイナイトのラス界面に生成した炭化物の円相当径を100nm以下と微細化することに加え、ラス界面に占める炭化物の量(以下、炭化物被覆率と言う)を30%以下とすることにより、拡散性水素の炭化物への集積を効果的に抑制し、より耐遅れ破壊特性を向上させる事が可能となる。   In addition, fine dispersion of carbide is extremely important for improving delayed fracture resistance. Most preferably, in addition to reducing the equivalent circle diameter of the carbide formed at the lath interface of martensite and lower bainite to 100 nm or less, the amount of carbide in the lath interface (hereinafter referred to as carbide coverage) is 30% or less. Thus, accumulation of diffusible hydrogen in the carbide can be effectively suppressed, and the delayed fracture resistance can be further improved.

マルテンサイトのラス界面に生成した炭化物の円相当径と炭化物被覆率の測定方法については、ナイタール等の適切な腐食液でエッチングした後に、走査型電子顕微鏡で組織写真を10視野以上撮影し、その写真を用いて求める。炭化物の円相当径は、例えば50個以上のラスの界面に生成した炭化物の大きさを画像解析により解析し、円相当径に換算する。また、炭化物被覆率は、50個以上のラス界面に生成した炭化物のラス界面に沿った長さと、ラスの界面長さを画像解析により計測し、炭化物のラス界面に沿った長さの総和を、ラス界面長さの総和で除し、100を掛けた数値とする。   About the measurement method of the equivalent circle diameter and carbide coverage of the carbide formed at the lath interface of martensite, after etching with an appropriate corrosive solution such as nital, the structure photograph was taken with a scanning electron microscope over 10 fields of view. Find using photos. For example, the equivalent circle diameter of carbide is converted into an equivalent circle diameter by analyzing the size of carbide generated at the interface of 50 or more laths by image analysis. In addition, the carbide coverage is measured by image analysis of the length along the lath interface of carbide generated at 50 or more lath interfaces, and the total length of the length along the lath interface of carbide. Divide by the sum of the lath interface lengths and multiply by 100.

3.製造条件について
以下に本発明の製造方法について説明する。
3. About manufacturing conditions The manufacturing method of this invention is demonstrated below.

なお本発明は、上述した組成を有する鋼を、転炉、電気炉等の溶製手段で溶製し、連続鋳造法または造塊〜分塊法等で常法によりスラブ等の鋼素材とすることができるが、鋼の溶製方法や鋳造方法を特定するものではない。   In the present invention, the steel having the above-described composition is melted by a melting means such as a converter or an electric furnace, and is made into a steel material such as a slab by a conventional method such as a continuous casting method or an ingot-bundling method. However, it does not specify a method for melting or casting steel.

圧延条件について
上述した組成を有する鋼片を、加熱炉でAc変態点以上に加熱する。加熱炉への鋼片の装入方法としては、鋳片をAr変態点以下に冷却することなく加熱炉に装入する熱片装入法や、一度冷却した鋳片を加熱炉に装入し、Ac変態点以上に再加熱する冷片装入法があるが、本発明ではいずれの方法も用いることができる。
The steel slab having the composition described above for rolling conditions, heated to Ac 3 transformation point or higher in a heating furnace. As a method of charging a steel slab into the heating furnace, a hot piece charging method in which the slab is charged into the heating furnace without being cooled below the Ar 3 transformation point, or a slab once cooled is charged into the heating furnace. However, although there is a cold piece charging method in which reheating is performed to the Ac 3 transformation point or higher, any method can be used in the present invention.

加熱炉でAc変態点以上に加熱するのは、鋼をオーステナイト組織一相に均一化するためであり、加熱温度としては、1100℃以上1250℃以下とするのが好ましい。特に靭性を重視する場合は1100℃以上1200℃以下とするのがより好ましい。 The reason why the steel is heated to the Ac 3 transformation point or higher in the heating furnace is to make the steel uniform in one phase of the austenite structure, and the heating temperature is preferably 1100 ° C. or higher and 1250 ° C. or lower. In particular, when emphasizing toughness, the temperature is more preferably 1100 ° C. or more and 1200 ° C. or less.

熱間圧延は、未再結晶温度域での累積圧下率を80%以下とし、Ar変態点以上で熱間圧延を終了し、続く加速冷却開始温度がAr変態点以上となるようにする。 Hot rolling, the cumulative reduction rate in the pre-recrystallization temperature region is 80% or less, and finished hot rolled at Ar 3 transformation point or higher, the accelerated cooling start temperature is made to be than the Ar 3 transformation point followed by .

累積圧下率は80%以下とするが、好ましい圧下率の範囲は、10〜80%の範囲である。なお、未再結晶温度域は圧延中に再結晶が起こらない温度域であり、本発明の鋼では930℃以下である。また、累積圧下率は(元厚−仕上厚)/元厚×100%で表される。   The cumulative rolling reduction is 80% or less, but a preferable rolling reduction range is 10 to 80%. The non-recrystallization temperature range is a temperature range where recrystallization does not occur during rolling, and is 930 ° C. or less in the steel of the present invention. The cumulative rolling reduction is expressed by (original thickness−finished thickness) / original thickness × 100%.

なお、Ar変態点は、下記式(1)により計算される値を用いる。 Incidentally, Ar 3 transformation point, a value that is calculated by the following equation (1).

Ac変態点は、下記式(2)により計算される値を用いる。 As the Ac 3 transformation point, a value calculated by the following equation (2) is used.

また、本発明鋼のうちNbを添加した鋼については、前記熱間圧延の未再結晶域圧下率を40%以上とし、かつ、熱間圧延終了温度を860℃以下とすると、焼戻し後に組織中に観察される炭化物の円相当径平均値が80nm以下となるため、さらに、耐遅れ破壊特性が向上する。   Moreover, about the steel which added Nb among this invention steel, when the unrecrystallized zone reduction of the said hot rolling shall be 40% or more, and the hot rolling completion temperature shall be 860 degrees C or less, in a structure | tissue after tempering Since the average equivalent circle diameter of the carbides observed in 1 is 80 nm or less, the delayed fracture resistance is further improved.

圧延後の冷却条件
熱間圧延終了後、母材強度および靭性を確保するため、Ar変態点以上の温度から250℃以下まで強制冷却を行う必要がある。
Cooling conditions after rolling In order to ensure the strength of the base metal and the toughness after the hot rolling is completed, it is necessary to perform forced cooling from a temperature equal to or higher than the Ar 3 transformation point to 250 ° C. or lower.

冷却停止温度が250℃以下になるまで冷却する理由は、オーステナイトからマルテンサイトへの変態を完了させ、母材を強化するためである。   The reason for cooling until the cooling stop temperature is 250 ° C. or lower is to complete the transformation from austenite to martensite and strengthen the base material.

本発明では、強度と靭性を両立するためにマルテンサイト+下部ベイナイト組織分率を95%以上とする。   In the present invention, in order to achieve both strength and toughness, the martensite + lower bainite structure fraction is set to 95% or more.

強制冷却時の冷却速度は、10℃/s以上とする。10℃/s未満では冷却時に、部分的にフェライト、パーライトが生成し易くなり、所望の強度、靭性を安定的に確保できないからである。   The cooling rate during forced cooling is 10 ° C./s or more. This is because if it is less than 10 ° C./s, ferrite and pearlite are likely to be partially formed during cooling, and the desired strength and toughness cannot be secured stably.

冷却方法は、直接焼入れ、加速冷却等の手法が用いられるが、冷却速度10℃/s以上、冷却停止温度250℃以下が得られれば冷却方法を特定するものではない。
なお、冷却速度は700〜500℃での平均冷却速度で規定する。この温度域がフェライトやパーライト等の軟質相が出易い温度領域であり、高強度を得るには、この温度領域を早く冷却する必要があるからである。
As a cooling method, techniques such as direct quenching and accelerated cooling are used, but the cooling method is not specified if a cooling rate of 10 ° C./s or more and a cooling stop temperature of 250 ° C. or less are obtained.
In addition, a cooling rate is prescribed | regulated by the average cooling rate in 700-500 degreeC. This is because this temperature region is a temperature region in which a soft phase such as ferrite or pearlite is likely to be produced, and it is necessary to cool this temperature region quickly in order to obtain high strength.

焼戻し条件
焼戻しは、圧延機および直接冷却もしくは加速冷却装置と同一の製造ライン上に直結して設置された加熱装置を用いて行うのが良い。これは、直結化により、圧延・冷却処理から焼戻し処理までに要する時間を短くすることが可能となり、生産性の向上、熱エネルギーの低減効果がもたらされるためである。
Tempering conditions Tempering is preferably performed using a heating device installed directly on the same production line as the rolling mill and the direct cooling or accelerated cooling device. This is because the time required from the rolling / cooling process to the tempering process can be shortened by the direct connection, thereby improving the productivity and reducing the heat energy.

焼戻しの加熱方式は、平均昇温速度が達成でき、加熱温度の上限・下限を管理できる方式であれば、誘導加熱、通電加熱、赤外線輻射加熱、雰囲気加熱等いずれを用いてもよい。   As the heating method for tempering, any method such as induction heating, current heating, infrared radiation heating, atmosphere heating, or the like may be used as long as the average temperature rising rate can be achieved and the upper and lower limits of the heating temperature can be managed.

焼戻しの温度条件は、加熱時の平均昇温速度を1℃/s以上とし、加熱温度の上限を400℃とする。平均昇温速度を1℃/s以上とし、焼戻し温度を400℃以下とすることで、ラス間に生成する炭化物のサイズを効果的に微細化し、従来鋼に比べて強度および靭性に優れた鋼板を得ることが可能となる。   The temperature condition for tempering is such that the average rate of temperature rise during heating is 1 ° C./s or more, and the upper limit of the heating temperature is 400 ° C. A steel plate that has an average temperature increase rate of 1 ° C./s or more and a tempering temperature of 400 ° C. or less, effectively reducing the size of carbides generated between the laths, and is superior in strength and toughness compared to conventional steel. Can be obtained.

一方、加熱温度が100℃未満の場合、脱水素効果が不十分であり、耐遅れ破壊特性の効果が十分に得られないので、加熱下限温度を100℃以上とする。   On the other hand, when the heating temperature is less than 100 ° C., the dehydrogenation effect is insufficient, and the effect of delayed fracture resistance is not sufficiently obtained.

また、平均昇温速度は冷却後、再加熱温度(100〜400℃)までの再加熱に必要な昇温量を再加熱に要した時間で割った値である。   The average rate of temperature increase is a value obtained by dividing the amount of temperature increase required for reheating up to the reheating temperature (100 to 400 ° C.) by the time required for reheating after cooling.

焼戻し時の昇温過程は、所定の平均昇温速度が得られればよく、直線的な温度履歴を取っても、途中温度で滞留するような温度履歴を取っても構わない。   The temperature raising process at the time of tempering is not limited as long as a predetermined average temperature rising rate is obtained, and a linear temperature history may be taken or a temperature history that stays at an intermediate temperature may be taken.

なお、上記した本発明の温度は、特に明記しない限り、いずれも板厚中心部の温度であり、表面実測温度からの計算により管理される。   The above-described temperatures of the present invention are all temperatures at the center of the plate thickness unless otherwise specified, and are managed by calculation from the surface measured temperature.

焼戻し温度における保持時間は、生産性・製造費用や析出物の粗大化に起因する靭性の劣化を防止するために、60s以下とするのが望ましい。   The holding time at the tempering temperature is preferably 60 s or less in order to prevent deterioration of toughness due to productivity / manufacturing cost and coarsening of precipitates.

焼戻し後の冷却速度については、冷却中の析出物の粗大化に起因する靭性の劣化を防止すべく、100℃以下までにおける板厚中心部の平均冷却速度を0.05℃/s以上、20℃/s以下とすることが望ましい。   With respect to the cooling rate after tempering, the average cooling rate at the central portion of the plate thickness up to 100 ° C. or lower is 0.05 ° C./s or higher, 20 to prevent deterioration of toughness due to coarsening of precipitates during cooling. It is desirable that the temperature is not more than ° C / s.

表1に示す化学成分の鋼種A〜Xの24種類を溶製してスラブを鋳造し、加熱炉で加熱後、圧延を行い鋼板とした。圧延後、引き続き直接焼入れし、次いで、ソレノイド型誘導加熱装置を用いて焼戻し処理を行った。   24 types of steel types A to X of chemical components shown in Table 1 were melted to cast a slab, heated in a heating furnace, and then rolled to obtain a steel plate. After rolling, the steel was then directly quenched and then tempered using a solenoid type induction heating device.

表2に鋼板製造条件、および得られた鋼板の降伏強度、引張強度、−40℃における吸収エネルギー(vE−40)、耐遅れ破壊安全度指数を示す。   Table 2 shows the steel sheet production conditions and the yield strength, tensile strength, absorbed energy at -40 ° C. (vE-40), and delayed fracture resistance safety index.

板厚中心部の平均昇温速度は鋼板の通板速度によって管理した。なお保持する場合には、鋼板を往復させて加熱することによって、±5℃の範囲で保持した。   The average heating rate at the center of the plate thickness was controlled by the plate passing rate of the steel plate. In addition, when hold | maintaining, it hold | maintained in the range of +/- 5 degreeC by reciprocating and heating a steel plate.

また、加熱後の冷却は空冷とした。焼戻し温度や焼入れ温度などの板厚中心部における温度は、放射温度計による表面の逐次における温度測定結果から、伝熱計算によって求めた。   The cooling after heating was air cooling. The temperature at the center of the plate thickness, such as the tempering temperature and the quenching temperature, was obtained by heat transfer calculation from the results of temperature measurement at the surface in succession by a radiation thermometer.

引張試験はJIS Z 2241に準拠して行い、板厚20mm以下ではJIS5号試験片により、板厚20mm超では板厚の1/4部から採取したJIS4号試験片により降伏強度および引張強度を測定した。   Tensile tests are performed in accordance with JIS Z 2241. Yield strength and tensile strength are measured with a JIS No. 5 test piece when the plate thickness is 20 mm or less, and with a JIS No. 4 test piece taken from 1/4 of the plate thickness when the plate thickness exceeds 20 mm. did.

靭性はJIS Z 2242に規定の衝撃試験片を採取し、板厚の1/4部より採取した試験片を用いたシャルピー衝撃試験によって得られる−40℃における吸収エネルギー(vE−40)で評価した。   Toughness was evaluated by absorbing energy (−40 ° C.) at −40 ° C. obtained by a Charpy impact test using a test specimen taken from ¼ part of the plate thickness after collecting the impact test specimen specified in JIS Z 2242. .

耐遅れ破壊安全度指数は、平滑丸棒試験片を用いて陰極水素チャージ法により試験片中の拡散性水素量が約0.5massppmになるように水素をチャージ後、試験片表面に亜鉛めっきを施すことにより水素を封入し、その後、1×10−6/sの歪速度にて引張試験を行い、破断した試験片の絞りを求め、一方、同様の歪速度にて水素チャージを行わない試験片の引張試験も行い、下記の式に従って評価した。
耐遅れ破壊安全度指数(%)=100×(X1/X0)
ここで、X0:実質的に拡散性水素を含まない試験片の絞り(%)
X1:拡散性水素を含む試験片の絞り(%)
各特性の目標値は、降伏応力が685MPa以上、引張強度が780MPa以上、vE−40が40J以上、耐遅れ破壊安全度指数の目標は80%以上とした。
The delayed fracture resistance index is obtained by charging hydrogen so that the amount of diffusible hydrogen in the test piece is about 0.5 massppm by the cathodic hydrogen charging method using a smooth round bar test piece, and then galvanizing the surface of the test piece. Hydrogen is sealed by applying, and then a tensile test is performed at a strain rate of 1 × 10 −6 / s to obtain a squeeze of the fractured test piece, while no hydrogen charge is performed at the same strain rate. A tensile test of the piece was also performed and evaluated according to the following formula.
Delayed fracture resistance index (%) = 100 x (X1 / X0)
Here, X0: Drawing of test specimen substantially not containing diffusible hydrogen (%)
X1: Test specimen containing diffusible hydrogen (%)
The target values for each characteristic were a yield stress of 685 MPa or more, a tensile strength of 780 MPa or more, a vE-40 of 40 J or more, and a delayed fracture resistance index target of 80% or more.

鋼板の組織の定量は、板厚1/4部付近についてナイタール腐食液で組織を現出し、光学顕微鏡で5視野観察を行い、画像解析による測定値の平均値で評価した。   Quantification of the structure of the steel plate was evaluated with an average value of measured values obtained by image analysis by revealing the structure with a nital etchant in the vicinity of ¼ part of the plate thickness, observing 5 fields of view with an optical microscope.

炭化物の円相当径および被覆率の測定は、板厚1/4部付近についてナイタール腐食液で組織を現出し、走査型電子顕微鏡で10視野観察し、画像解析による測定値の平均で評価した。   For the measurement of the equivalent circle diameter and the coverage of the carbide, the structure was revealed with a nital etchant in the vicinity of ¼ part of the plate thickness, 10 fields of view were observed with a scanning electron microscope, and the average of the measured values by image analysis was evaluated.

ただし、鋼板の強度が得られなかった、No.18、25、28、30、31については炭化物の円相当径および被覆率は測定しなかった。   However, the strength of the steel sheet was not obtained, No. For 18, 25, 28, 30, and 31, the equivalent circle diameter and coverage of the carbide were not measured.

表2に示すように、本発明法により製造した発明例No.1〜No.16は化学成分、製造方法、が本発明の範囲内であり、良好な強度と靭性および耐遅れ破壊特性を有する鋼板を得ることができる。   As shown in Table 2, Invention Example No. manufactured by the method of the present invention was used. 1-No. No. 16 has chemical components and production methods within the scope of the present invention, and a steel sheet having good strength and toughness and delayed fracture resistance can be obtained.

また、Nb添加した発明例のうち圧延仕上温度が870℃以下で、かつ、未再結晶域圧下率が40%以上で製造した鋼は炭化物の平均円相当径が80nm以下となり、さらに耐遅れ破壊特性が優れている。   Further, among the Nb-added invention examples, the steel manufactured with a rolling finishing temperature of 870 ° C. or less and an unrecrystallized zone reduction of 40% or more has an average equivalent circle diameter of carbide of 80 nm or less, and is further resistant to delayed fracture. Excellent characteristics.

これに対して、成分が本発明から外れる比較例No.17〜No.24は、強度、靭性および耐遅れ破壊特性のいずれか1つ以上の特性が目標値に達していない。また、製造条件が本発明から外れる比較例No.25〜No.33の鋼板は、強度、靭性および耐遅れ破壊特性のいずれか1つ以上の特性が目標値に達していない。   On the other hand, comparative example No. in which a component remove | deviates from this invention. 17-No. In No. 24, any one or more of strength, toughness, and delayed fracture resistance does not reach the target value. In addition, Comparative Example No. in which the manufacturing conditions deviate from the present invention. 25-No. As for 33 steel plate, any one or more of strength, toughness and delayed fracture resistance does not reach the target value.

Claims (6)

質量%で、C:0.02〜0.25%、Si:0.01〜0.8%、Mn:0.5〜2%、P:0.010%以下、S:0.003%以下、Al:0.005〜0.1%、N:0.0005〜0.008%を含有し、残部がFeおよび不可避的不純物からなる鋼をAc変態点以上に加熱し、未再結晶温度域での累積圧下率を80%以下とする熱間圧延を行い、Ar変態点以上で熱間圧延を終了し、引き続きAr変態点以上から10℃/s以上の冷却速度で250℃以下の温度まで冷却後、5℃/s以上の平均昇温速度で再加熱し、最高到達温度を100〜400℃の範囲とし、前記最高到達温度における保持時間を60s以内とする焼戻し処理を行うことを特徴とする、降伏応力が685MPa以上、引張強度が780MPa以上である耐遅れ破壊特性に優れた高張力鋼板の製造方法。 In mass%, C: 0.02-0.25%, Si: 0.01-0.8%, Mn: 0.5-2%, P: 0.010% or less, S: 0.003% or less , Al: 0.005 to 0.1% N: containing 0.0005 to 0.008%, heating the steel and the balance being Fe and unavoidable impurities than Ac 3 transformation point, non-recrystallized temperature subjected to hot rolling to a cumulative reduction rate at frequency 80% or less, Ar 3 to exit the hot rolling at lower than the transformation point, continuing Ar 3 below 250 ° C. at 10 ° C. / s or more from the above transformation point of cooling rate After cooling to a temperature of 5 ° C / s, reheating at an average temperature increase rate of 5 ° C / s or more, setting the maximum temperature to 100 to 400 ° C, and performing a tempering process with the holding time at the maximum temperature being within 60s. The yield stress is 685 MPa or more, and the tensile strength is 780 MPa or more. Delayed production method excellent high-tensile steel plate in fracture characteristics that. 質量%で、C:0.02〜0.25%、Si:0.01〜0.8%、Mn:0.5〜1.5%、P:0.010%以下、S:0.003%以下、Al:0.005〜0.1%、N:0.0005〜0.008%を含有し、更に、Mo:0.01〜1%、Nb:0.001〜0.1%、V:0.001〜0.5%、Ti:0.001〜0.1%、Cu:2%以下、Ni:4%以下、Cr:2%以下、W:2%以下の中から選ばれる一種以上を含有し、残部がFeおよび不可避的不純物からなる鋼をAc 変態点以上に加熱し、未再結晶温度域での累積圧下率を80%以下とする熱間圧延を行い、Ar 変態点以上で熱間圧延を終了し、引き続きAr 変態点以上から10℃/s以上の冷却速度で250℃以下の温度まで冷却後、5℃/s以上の平均昇温速度で再加熱し、最高到達温度を100〜400℃の範囲とし、前記最高到達温度における保持時間を60s以内とする焼戻し処理を行うことを特徴とする、降伏応力が685MPa以上、引張強度が780MPa以上である耐遅れ破壊特性に優れた高張力鋼板の製造方法。 In mass%, C: 0.02 to 0.25%, Si: 0.01 to 0.8%, Mn: 0.5 to 1.5%, P: 0.010% or less, S: 0.003 % or less, Al: 0.005 to 0.1%, N: containing from 0.0005 to 0.008%, further, M o: 0.01~1%, Nb : 0.001~0.1% V: 0.001 to 0.5%, Ti: 0.001 to 0.1%, Cu: 2% or less, Ni: 4% or less, Cr: 2% or less, W: 2% or less A steel containing the balance of Fe and inevitable impurities is heated to the Ac 3 transformation point or higher, and hot rolling is performed so that the cumulative reduction in the non-recrystallization temperature range is 80% or less, and Ar 3 Exit hot rolling at lower than the transformation point, subsequently cooled to a temperature of 250 ° C. or less than the Ar 3 transformation point from 10 ° C. / s or more cooling rate, 5 ° C. / s or higher Reheating at an average rate of temperature increase of 100 to 400 ° C., and performing a tempering treatment with a holding time at the maximum temperature of 60 s or less . The yield stress is 685 MPa or more, method for producing a high tensile steel sheet tensile strength and excellent resistance to delayed fracture resistance Ru der least 780 MPa. 前記鋼に、更に、質量%で、B:0.0003%〜0.003%、Ca:0.01%以下、REM:0.02%以下の中から選ばれる一種以上を含有することを特徴とする、降伏応力が685MPa以上、引張強度が780MPa以上である請求項1または2に記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   The steel further contains one or more selected from B: 0.0003% to 0.003%, Ca: 0.01% or less, and REM: 0.02% or less in mass%. The method for producing a high-tensile steel sheet having excellent delayed fracture resistance according to claim 1 or 2, wherein the yield stress is 685 MPa or more and the tensile strength is 780 MPa or more. 圧延後の再加熱を圧延機および冷却装置と同一の製造ライン上に設置された加熱装置を用いて行うことを特徴とする、降伏応力が685MPa以上、引張強度が780MPa以上である請求項1乃至3の何れかに記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   The yield stress is 685 MPa or more and the tensile strength is 780 MPa or more, wherein reheating after rolling is performed using a heating device installed on the same production line as the rolling mill and the cooling device. 4. A method for producing a high-tensile steel sheet excellent in delayed fracture resistance according to any one of 3 above. 前記鋼が、質量%で、Nb:0.001〜0.1%を含有し、未再結晶温度域での累積圧下率を40%以上、熱間圧延終了温度を860℃以下とすることを特徴とする、降伏応力が685MPa以上、引張強度が780MPa以上である請求項1乃至4の何れかに記載の耐遅れ破壊特性に優れた高張力鋼板の製造方法。   The steel contains, by mass%, Nb: 0.001 to 0.1%, a cumulative rolling reduction in the non-recrystallization temperature range of 40% or more, and a hot rolling end temperature of 860 ° C. or less. The method for producing a high-tensile steel sheet having excellent delayed fracture resistance according to any one of claims 1 to 4, wherein the yield stress is 685 MPa or more and the tensile strength is 780 MPa or more. 前記高張力鋼板の金属組織はマルテンサイトおよび下部ベイナイトの組織分率を合計で95面積%以上とし、前記金属組織中のラス界面における炭化物の円相当径は100nm以下、炭化物被覆率は30%以下であることを特徴とする請求項1乃至5の何れかに記載の、降伏応力が685MPa以上、引張強度が780MPa以上である耐遅れ破壊特性に優れた高張力鋼板の製造方法。   The metal structure of the high-tensile steel sheet has a total structure fraction of martensite and lower bainite of 95 area% or more, the equivalent circle diameter of carbide at the lath interface in the metal structure is 100 nm or less, and the carbide coverage is 30% or less. The method for producing a high-strength steel sheet having excellent delayed fracture resistance, wherein the yield stress is 685 MPa or more and the tensile strength is 780 MPa or more, according to any one of claims 1 to 5.
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Family Cites Families (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
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JP5157066B2 (en) * 2004-12-28 2013-03-06 Jfeスチール株式会社 A method for producing a high-strength, high-toughness thick steel plate excellent in cut cracking resistance and DWTT characteristics.
JP5157072B2 (en) * 2005-03-29 2013-03-06 Jfeスチール株式会社 Manufacturing method of high strength and high toughness thick steel plate with excellent tensile strength of 900 MPa and excellent in cutting crack resistance
JP5277672B2 (en) * 2007-03-29 2013-08-28 Jfeスチール株式会社 High strength steel plate with excellent delayed fracture resistance and method for producing the same
AU2009294126B2 (en) * 2008-09-17 2011-03-10 Nippon Steel Corporation High-strength steel plate and producing method thereof
JP5354164B2 (en) * 2008-12-09 2013-11-27 Jfeスチール株式会社 Low yield ratio high strength thick steel plate and method for producing the same

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