JP4445161B2 - Manufacturing method of thick steel plate with excellent fatigue strength - Google Patents

Description

【００６６】
表１の化学組成の鋼片を、表２に示す条件の熱間圧延、表３に示す条件の熱処理、さらに一部は焼戻しを施して、板厚２５ｍｍ又は５０ｍｍの鋼板に製造し、室温の引張特性、２ｍｍＶノッチシャルピー衝撃特性、さらに溶接継手の疲労特性を調査した。引張試験片及びシャルピー衝撃試験片は板厚中心部から圧延方向に直角（Ｃ方向）に採取した。引張特性は室温で測定し、シャルピー衝撃特性は５０％破面遷移温度（ｖＴｒｓ）で評価した。疲労試験は、構造物の溶接止端部から疲労き裂が発生し、母材部を伝播する場合の疲労特性を評価するために、図１に示す廻し溶接継手について行った。試験片は、鋼板から鋼板長手方向長さ：３００ｍｍ、幅方向長さ：８０ｍｍ、板厚：２５ｍｍ（２５ｍｍ厚材については全厚、５０ｍｍ厚材については表面から採取）、のサイズで試験板を採取し、幅：１０ｍｍ、長さ：３０ｍｍ、高さ：３０ｍｍのリブ板を炭酸ガス溶接（Ｃｏ２溶接）により、試験板の中央に廻し溶接で溶接した。この際の炭酸ガス溶接は、化学組成が、Ｃ：０．０６ｍａｓｓ％、Ｓｉ：０．５ｍａｓｓ％、Ｍｎ：１．４ｍａｓｓ％、である１．４ｍｍ径の溶接ワイヤを用いて、電流：２７０Ａ、電圧：３０Ｖ、溶接速度：２０ｃｍ／ｍｉｎ．で行った。疲労試験は、荷重支点Ｐのスパンを、下スパン：７０ｍｍ、上スパン：２２０ｍｍとして、最大荷重（Ｐｍａｘ）：５５００ｋｇｆで応力比（Ｒ）：０．１の繰り返し応力負荷を加え、疲労寿命を測定した。
【００６７】
【表２】

【００６８】
【表３】

【００６９】
表２、３における注）は以下の通りである。
注１） 本発明で必須の、「８５０℃以下、終了温度Ａｒ_３以上、累積圧下率３０％以上の圧延」に先立つ圧延
注２） 本発明で必須の「８５０℃以下、終了温度Ａｒ_３以上、累積圧下率３０％以上の圧延」及び比較例において本圧延に対応する条件での圧延
注３） 圧延後、空冷での実測値
注４） 手段は空冷（ＡＣ）または加速冷却（ＡｃＣ）。空冷の場合は全て室温まで冷却。冷却速度は圧延終了から５００℃までの平均。
注５） 加熱温度から冷却停止温度までの平均。
注６） 焼戻しの保持は、３０分〜２時間の範囲。冷却は全て空冷。
【００７０】
鋼板の組織形態（フェライト粒径、硬質第二相の種類、割合、ビッカース硬さ、アスペクト比）と機械的性質を表４に示す。なお、組織の定量は、鋼板長手方向に平行な板厚断面において、表面下１ｍｍ、板厚の１／４、板厚中心部の光学顕微鏡組織を撮影し、各位置での平均フェライト粒径、硬質第二相の割合、アスペクト比を求め、３観察位置の平均値を示した。硬質第二相の硬さは、表面下１ｍｍ、板厚の１／４、板厚中心部、各位置で荷重５〜１０ｇのマイクロビッカース硬さを１０点以上測定し、各位置での平均値をさらに平均した。
【００７１】
【表４】

【００７２】
表２〜４の内の鋼板番号Ａ１〜Ａ１３は、本発明の化学組成と組織に関する要件を全て満足している鋼板であり、いずれも構造用鋼として必要な強度、靭性（２ｍｍＶノッチシャルピー衝撃特性）を有しているだけでなく、良好な継手疲労特性も有していることが明らかである。
【００７３】
一方、鋼板番号Ｂ１〜Ｂ１１は、本発明のいずれかの要件を満足していない、比較の鋼板であり、本発明の鋼板に比べて、継手疲労特性や靭性が劣っていることが明白である。
【００７４】
鋼板番号Ｂ１〜Ｂ５は、化学組成が本発明を満足していないために、本発明の組織要件を満足できないか、あるいは本発明の組織要件を満足しているにも関わらず、良好な特性を達成できなかった例である。
【００７５】
すなわち、鋼板番号Ｂ１は、Ｃ量が過少ないため、硬質第二相の焼入性が低く、全面的にパーライトとベイナイトとの混合組織となり、硬質第二相の硬さが本発明の下限よりも低くなっている。そのため、本発明の鋼板に比べて継手疲労特性が大きく劣っている。
【００７６】
鋼板番号Ｂ２は、逆にＣ量が過剰な比較例であり、硬質第二相の硬さが過剰で非常に脆いため、靭性が極めて劣位である。継手疲労特性は比較的良好であるが、疲労試験においても硬質相が脆性破壊する影響で、本発明に比べて、継手疲労特性は若干劣る。
【００７７】
鋼板番号Ｂ３、Ｂ４は、各々Ｍｎ量、Ｐ量がが過剰なため、特に靭性の劣化が著しい。また、靭性の劣化に起因して、継手疲労試験においても、フェライト相、硬質相での脆性破壊が一部生じるために、本発明に比べて、継手疲労特性は劣る。
【００７８】
鋼板番号Ｂ５も、Ｓ量が過剰なため、靭性、継手疲労特性が本発明に比べて劣る。特に継手疲労特性の劣化が大きい。
【００７９】
鋼板番号Ｂ６〜Ｂ１１は、化学組成は本発明を満足しているものの、組織要件が本発明を満足していないために、継手疲労特性が劣っているか、靭性が構造用鋼として十分でない例である。
【００８０】
すなわち、鋼板番号Ｂ６は、最終の熱処理の加熱温度がＡＣ₃変態点超と、本発明の製造方法の上限を超えているため、フェライト相が残存せず、加速冷却後の組織が全面ベイナイトとマルテンサイトの混合組織となり、フェライトと硬質第二相との二相組織とはなっていないため、継手疲労特性の改善が全く認められない。
【００８１】
一方、鋼板番号Ｂ７は、最終の熱処理の加熱温度が本発明の製造方法の下限未満のため、加熱段階において硬質相形成に必要な一定量のオーステナイト化がなされず、その結果、最終組織中に、本発明で必須のベイナイト、マルテンサイトから構成され、硬さが２３０以上の硬質第二相が形成されず、継手疲労特性が本発明に比べて極めて低い。
【００８２】
また、鋼板番号Ｂ８は、全圧下比が小さいために、硬質第二相のアスペクト比が過小で、本発明を満足しておらず、本発明に比べて継手疲労特性が劣る。
【００８３】
鋼板番号Ｂ９は、最終の熱処理の冷却速度が小さいため、硬質第二相の硬さが過小で、き裂伝播抑制に有効に機能せず、従って、継手疲労特性が顕著に劣る。
【００８４】
鋼板番号Ｂ１０は、熱間圧延において、８５０℃以下から開始する圧延行っていないため、フェライト粒界が本発明の範囲をはずれて過大であるため、靭性が劣り、また、硬質第二相の層状化が助長されないため、同程度の全圧下比で圧延された本発明鋼に比べて、硬質第二相のアスペクト比が本発明の範囲内ではあるが、わずかに小さくなっており、継手疲労特性は、比較例に比べれば十分優れているが、他の組織要件が類似の本発明鋼に比べてわずかながら劣る。
【００８５】
鋼板番号Ｂ１１は熱間圧延後に冷却速度の大きい加速冷却を施したために、圧延終了段階での第二相の層状化がほとんどなされておらず、そのため、最終熱処理後の硬質第二相のアスペクト比が過小となって、継手疲労特性が改善されていない。
【００８６】
以上の実施例から、本発明によれば、構造用鋼として十分高い靭性を確保しながら、優れた継手疲労特性を得ることが可能であることが明白である。
【００８７】
【発明の効果】
本発明は疲労強度が必要とされる溶接構造部材に用いられる、引張強さが４００ＭＰａ級以上の厚鋼板において、従来、溶接部では向上が困難とされてきた、継手疲労特性の向上を特殊な合金元素や複雑な製造プロセスに頼ることなく製造できる点で、産業上の有用性は極めて大きい。
【図面の簡単な説明】
【図１】溶接継手疲労特性評価のための、廻し溶接４点曲げ試験方法の概念図である。
【符号の説明】
Ｐ：荷重支点[0001]
BACKGROUND OF THE INVENTION
  The present invention is a thick steel having a tensile strength of 400 MPa or more used for welded structural members that require fatigue strength.PlankIt relates to a manufacturing method. The steel sheet of the present invention can be used in general for welded steel structures such as marine structures, pressure vessels, shipbuilding, bridges, buildings, line pipes, etc., but especially marine structures, shipbuilding, bridges that require fatigue strength. It is useful as a structural steel plate for construction structures. In addition, the present invention can also be applied to steel pipe materials or shape steels that are used as structural members and require fatigue strength.
[0002]
[Prior art]
With the increase in the size of welded structures and the demand for environmental protection, there has been a demand for increased reliability of structural members. The current structure is generally a welded structure, and the failure modes assumed for welded structures include fatigue failure, brittle failure, ductile failure, etc. Of these, the most frequent failure modes are: Brittle fracture or fatigue fracture from initial defects, and further brittle fracture following fatigue fracture. In addition, these forms of destruction are difficult to prevent only by considering the design of the structure, and often cause sudden collapse of the structure. From the viewpoint of ensuring the safety of the structure, it is also possible to prevent it. Is the most required form of destruction.
[0003]
For brittle fracture, there are means for improving the chemical composition such as addition of Ni and optimization of the transformation structure, and the manufacturing method can be improved by refinement of the structure by controlled rolling or thermomechanical treatment. On the other hand, in the case of fatigue characteristics, it is possible to improve the smooth member by improving the strength. It was considered impossible to improve fatigue strength (joint fatigue strength) by metallurgical means. That is, in a structure where fatigue strength is a problem, the design strength cannot be increased even when high strength steel is used, and the advantage of using high strength steel cannot be obtained. Therefore, conventionally, in such a welded structure, joint fatigue strength has been improved by so-called toe treatment for improving the shape of the weld toe part that is a stress concentration part. For example, a method of cutting the toe by a grinder to increase the toe radius, a method of remelting the toe part by TIG welding to make the toe shape smooth (for example, Japanese Patent Publication No. 54-30386), shot peening In other words, a compressive stress is generated in the toe portion.
[0004]
However, since these toe treatments are very time-consuming, a means for improving the joint fatigue strength of the steel material itself has been awaited in order to reduce costs and improve productivity.
[0005]
Recently, several steel materials having good joint fatigue strength have been proposed in response to such demands. For example, a technique (Japanese Patent Laid-Open No. 8-73983) that can improve the fatigue strength of HAZ by setting the structure of the weld heat affected zone (HAZ) to ferrite (α) is shown. However, since the present technology requires that the HAZ structure be a ferrite structure, there is a limit to the strength level of the steel material that can be manufactured, and it is not possible to manufacture a high-strength steel material having a tensile strength exceeding 780 MPa.
[0006]
Several means for improving the joint fatigue strength of high-strength steel with a tensile strength of 590 MPa or more have been proposed. For improving the fatigue crack initiation and propagation characteristics of the HAZ bainite structure, high Si (Japanese Patent Laid-Open No. Hei 8-209295) is proposed. There are reports that high Nb (Japanese Patent Laid-Open No. 10-1743) is effective. However, when both Si and Nb are added in a large amount, it is an element that greatly deteriorates toughness, and there is also a concern that manufacturing problems such as cracking of a steel piece may occur.
[0007]
Each of the above prior arts is a means for improving the generation of fatigue cracks in the HAZ structure and the propagation of fatigue cracks in the HAZ. However, since HAZ is greatly affected by the stress concentration at the toe, depending on the shape of the toe The effect may not occur or may be small.
[0008]
In order to improve the joint fatigue strength regardless of the shape of the toe, it is effective to delay the propagation of the fatigue crack generated from the toe at the base material. Based on this concept, a technology for improving the fatigue crack growth characteristics of a base material by forming a base material structure in which coarse ferrite is dispersed in a fine grain structure with an average ferrite grain size of 20 μm or less (special feature). (Kaihei 7-90481) is disclosed. However, in this case as well, only a steel material having a tensile strength of about 580 MPa class can be manufactured because of the necessity for a ferrite main structure.
[0009]
Furthermore, as a technique to increase fatigue strength by suppressing fatigue crack propagation in the base metal, in a structure consisting of ferrite and a hard second phase, there is a certain relationship between the hardness of the ferrite and the hardness of the hard second phase. Japanese Patent Application Laid-Open No. 11-1742 discloses a technique for defining the second phase form (aspect ratio, interval) or / and texture. This technology is one of the best methods for suppressing fatigue crack propagation among the currently shown technologies. However, the cumulative reduction in the two-phase region to the ferrite region is necessary for the formation of microstructure and texture. Since it is necessary to increase the rate, there are problems such as deterioration of productivity and deterioration of the steel plate shape.
[0010]
[Problems to be solved by the invention]
  The present invention relates to a thick steel plate having a tensile strength of 400 MPa or more used for welded structural members.Manufacturing methodIn order to improve the joint fatigue strength regardless of the shape of the toe, a thick steel plate with excellent fatigue crack propagation characteristics of the base metal is added in a large amount of special or expensive alloy elements and the productivity is inferior. Or without complicated manufacturing methodsA manufacturing method for thick steel plates with excellent fatigue strengthThe issue is to provide.
[0011]
[Means for Solving the Problems]
Means for improving joint fatigue strength without depending on the shape of the toe of the joint by improving the fatigue crack propagation characteristics of the base material in a steel material having a tensile strength of 400 MPa class or more. From the detailed experimental results of the relationship between fatigue crack growth behavior and steel microstructure. That is, when a fatigue crack generated from the stress concentration part of the joint toe portion propagates in the thickness direction, in the mixed structure of ferrite and the hard second phase having an appropriate form and characteristics, the interface between both structures or In many cases, fatigue crack stagnation, bending, branching, etc. occur in the vicinity of the interface, and when the crack propagates to the hard second phase, the progress of the fatigue crack in the hard second phase is remarkably suppressed. Is done. These comprehensive effects greatly reduce the macro fatigue crack propagation rate in the base metal, and in order to suppress such fatigue crack growth, the structure of the ferrite phase and the hard second phase It has been found that two-phase rolling is not necessarily an essential requirement if the form and properties are optimized.
[0012]
Furthermore, the present inventors have established, based on detailed experiments, industrially most preferable means for forming a microstructure suitable for the fatigue crack propagation characteristics of the base material.
[0013]
This invention is invented based on the above knowledge, and the summary is as follows.
[0017]
  (1)% By mass
C: 0.04-0.3%
Si: 0.01-2%
Mn: 0.1 to 3%
Al: 0.001 to 0.1%,
N: 0.001 to 0.01%
Containing
P: 0.02% or less,
S: 0.01% or less
The balance consists of iron and inevitable impuritiesSteel slab with components_ThreeReheated to transformation temperature to 1250 ° C, start temperature is 850 ° C or less, end temperature is Ar_ThreeAfter rolling at a transformation point or higher and including a rolling with a cumulative reduction ratio of 30% or more and a total rolling ratio of 5 or more, the steel is cooled to 500 ° C. or less at a cooling rate of 5 ° C./s or less, and then (AC₁Transformation point + 30 ° C) to (AC_ThreeTransformation point-50 ° C) and cooled to 400 ° C or less at 5-100 ° C / sThe ferrite and the hard second phase in the plate thickness cross-sectional structure parallel to the longitudinal direction of the steel sheet satisfy all the following conditions (1) to (4). The hard second phase structure is either bainite, martensite or a mixed structure of bothA method for producing a thick steel plate having excellent fatigue strength.
(1) Average ferrite particle size: 20 μm or less
(2) Ratio of hard second phase: 10 to 70%
(3) Average Vickers hardness of hard second phase: 230 or more
(4) Average aspect ratio of hard second phase (average length in the longitudinal direction of steel sheet / average length in the thickness direction): 10 or more
(2) Furthermore, the steel slab is mass%,
Ni: 0.01-6%,
Cu: 0.01 to 1.5%,
Cr: 0.01-2%
Mo: 0.01-2%
W: 0.01-2%
Ti: 0.003 to 0.1%,
V: 0.005-0.5%
Nb: 0.003 to 0.2%,
Zr: 0.003 to 0.1%,
Ta: 0.005 to 0.2%,
B: 0.0002 to 0.005%
The manufacturing method of the thick steel plate excellent in the fatigue strength as described in said (1) characterized by containing 1 type, or 2 or more types.
(3) Further, the steel slab is mass%,
Mg: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
REM: 0.005-0.1%
The manufacturing method of the thick steel plate excellent in the fatigue strength as described in said (1) or (2) characterized by containing 1 type, or 2 or more types.
[0018]
  (4) Furthermore,Tempering at 250 to 600 ° C.Any of 1) to (3)The manufacturing method of the thick steel plate excellent in the fatigue strength of description.
[0019]
DETAILED DESCRIPTION OF THE INVENTION
An object of the present invention is to secure joint fatigue strength by improving fatigue crack propagation characteristics of a base material. The most important requirement is that the base material structure is “having a structure composed of ferrite and a hard second phase, and in the plate thickness cross-sectional structure parallel to the longitudinal direction of the steel sheet, the ferrite and the hard second phase are 1) Average ferrite particle size: 20 μm or less, 2) Hard second phase ratio consisting of bainite, martensite, or a mixed structure of both: 10 to 70%, 3) Average Vickers hardness of hard second phase: 230 As described above, (4) the average aspect ratio of the hard second phase (average length in the longitudinal direction of the steel sheet / average length in the thickness direction): satisfies the condition of 10 or more.
[0020]
As a technology to increase fatigue strength by suppressing fatigue crack propagation in the base metal, in the structure consisting of ferrite and hard second phase, a certain relationship was defined between the hardness of ferrite and the hardness of hard second phase. In the above, a technique defining the form of the second phase (aspect ratio, spacing) and / or texture has already been disclosed in JP-A-11-1742. The present invention is also similar in that it suppresses fatigue crack propagation by the structure consisting of ferrite and hard second phase, but the relationship between the structure and properties of ferrite and second phase and fatigue crack propagation behavior is observed in detail. As a result of the study, it was newly found that the fatigue strength characteristics can be further improved by using ferrite and a hard second phase characterized by the above (1) to (4), which are different from the above techniques. I found it. In other words, in order to improve the fatigue crack propagation characteristics of the base material, bending and branching of the fatigue cracks and the frequent entry into the hard second phase having appropriate characteristics can be achieved. The fatigue crack growth rate in the two phases is preferable because it is extremely small, and for the reliable bending, branching and penetration of the fatigue crack into the hard second phase, It has been found that the aspect ratio of the two phases needs to be very large as 10 or more, and that the hardness of the hard second phase is more important than the ratio of the hardness of the ferrite and the hard second phase. It was.
[0021]
Based on the above basic knowledge, based on detailed experiments, it was specifically derived that it is necessary to satisfy the organizational requirements (1) to (4) at the same time. The following is a more detailed explanation for each organizational requirement. In the present invention, it is assumed that the structure is composed of ferrite and a hard second phase. This is because in order to cause bending and branching of a fatigue crack, a mixed structure of ferrite and hard second phase is essential regardless of the form and characteristics of the hard second phase. In this case, it is preferable that the hard second phase is only the hard second phase of the present invention. However, as long as the present invention is satisfied, other phases such as pearlite or pseudo-pearlite may be partially included. The effect of the present invention is not inhibited.
[0022]
(1) Average ferrite grain size of 20 μm or less: Although this requirement is not directly related to improvement of fatigue characteristics, in the structure containing the hard second phase of the present invention, the hard phase is basically brittle and adversely affects toughness. In order to secure toughness as a whole steel sheet, the average ferrite grain size is set to 20 μm or less. If the average ferrite grain size is 20 μm or less, good toughness of 0 ° C. or less can be achieved at the fracture surface transition temperature (vTrs) in the 2 mmV notch Charpy impact test. On the other hand, when the average ferrite particle size exceeds 20 μm, there is a concern that vTrs deteriorates to a room temperature or more, and there is also a possibility that brittle fracture occurs during fatigue and fatigue crack propagation characteristics deteriorate.
[0023]
(2) The ratio of hard second phase consisting of bainite or martensite or a mixed structure of both is 10 to 70%: First, the hard second phase is bainite or martensite, each single-phase structure, or a mixed structure of both. Is essential. The structure phase is a fine lath structure, and if the Vickers hardness is 230 or more, it exists in a fine and high density even if carbide is present. However, it is extremely effective for suppressing crack growth when the crack enters the almost hard second phase. In addition, the bainite or martensite phase that transforms at a low temperature generates residual stress near the interface with the adjacent ferrite phase during transformation, increasing the probability of bending and branching of fatigue cracks. On the other hand, for example, pearlite and pseudo-pearlite have a layered structure of ferrite and cementite in the micro, so cracks can be selected and propagated in soft ferrite between layers, and cementite is brittle. The fatigue crack growth delay effect is smaller than that of bainite or martensite because, in many cases, cracks are not generated due to the occurrence of cracks near the crack tip. In the present invention, the ratio of the hard second phase composed of bainite, martensite, or a mixed structure of both is further limited to 10 to 70%. When trying to suppress fatigue crack growth by the hard second phase, naturally, the greater the probability that a hard second phase exists at the crack leading edge, the more the crack bends, branches, and even the cracks harder. The frequency of progress in the two phases will also increase. In order for the fatigue crack delay effect due to the hard second phase to be clearly reflected in the final increase in fatigue life, at least 10% of the hard second phase is required. On the other hand, to a certain extent, the fatigue properties improve as the proportion of the hard second phase increases. However, if the amount of the hard second phase is excessive, the proportion of the ferrite that is the soft phase becomes extremely small and substantially ferrite. It is difficult to say a two-phase structure of a hard second phase, and conversely, bending and branching of a fatigue crack are difficult to occur. Further, the hard second phase is not preferable for toughness, and if it exists in excess, there is a concern that the toughness is greatly deteriorated even if the ferrite grain size is reduced to 20 μm or less. The upper limit of the ratio of the hard second phase for the above reasons was examined based on the experimental results, and was set to 70% in the present invention.
[0024]
(3) The average Vickers hardness of the hard second phase is 230 or more: cracks and branches at the ferrite / hard second phase interface or in the vicinity of the interface, and the crack growth rate in the hard second phase is reduced. In order to sufficiently exhibit the crack propagation delay effect, it is necessary to optimize the hardness and form of the hard second phase. As a result of examining the relationship between the hardness and form of the hard second phase in detail and the fatigue crack propagation behavior, first, regarding the hardness of the hard second phase, it is not the hardness difference or ratio with ferrite, but the hardness The conclusion was that the hardness of the second phase itself is the most important, and it is necessary that the average hardness measured by Vickers hardness is 230 or more. If the average Vickers hardness is 230 or more, the plastic deformation at the tip of the fatigue crack can be reliably restrained, and it works effectively for bending and branching of the crack. Also, fatigue crack delay in the hard second phase is ensured. If the average Vickers hardness is less than 230, these effects cannot be expected with certainty. In addition, when it is not necessary to consider toughness so much depending on the application, it is preferable that the average Vickers hardness is 300 or more as a condition that is more effective in fatigue characteristics. When the Vickers hardness of the hard second phase is 300 or more, the crack retarding effect in the hard phase is more exhibited. On the other hand, when placing importance on toughness, it is more preferable to set the upper limit of the average Vickers hardness to 800. Since the hard second phase having an average Vickers hardness exceeding 800 has extremely poor toughness, an adverse effect on toughness may be unavoidable. In addition, depending on the composition, the hard second phase tends to brittle fracture during the fatigue test, which is not preferable for fatigue characteristics.
[0025]
(4) The average aspect ratio of the hard second phase (average length in the longitudinal direction of the steel sheet / average length in the thickness direction) is 10 or more: Probability that the hard second phase is present at the tip of the developing fatigue crack, In order to make bending and branching more reliable and to increase the length of the branch crack and the detour distance, it is preferable to form the hard second phase in a band shape. Since the fatigue crack generated at the joint toe propagates in the plate thickness direction, it is preferable to form a band-like structure in the longitudinal direction of the steel plate. When the index of the form of the band structure is the average aspect ratio (average length of the second steel plate in the longitudinal direction / average length in the thickness direction) in the thickness cross-sectional structure parallel to the longitudinal direction of the steel plate, the above effect is obtained. In order to exhibit, the average aspect ratio needs to be 10 or more. If the average aspect ratio is less than 10, considering the inevitable variation of the microstructure in the steel sheet, there is a possibility that the crack can propagate to a place other than the hard second phase without making a large detour. Therefore, it is not preferable. In addition, regarding the form of the hard second phase, as long as the ratio and the aspect ratio are in the specified range, the effect is sufficiently exhibited even if the individual sizes are somewhat different. However, it is more preferable that the average thickness of the hard second phase be 1 μm or more in order to further effectively utilize the crack retardation effect in the hard phase.
[0026]
As described above, the structural requirements (1) to (4) are essential for improving the fatigue strength by delaying the fatigue crack of the base material, which is the object of the present invention. Need to be specifically limited. The reasons for limiting the chemical composition in the present invention will be described below.
[0027]
First, C is an effective component for increasing the hardness of the hard second phase. If it is less than 0.04%, it is not easy to stably present 10% or more of the hard second phase having a Vickers hardness of 230 or more. Therefore, in the present invention, the lower limit of C is 0.04%. However, excessive content exceeding 0.3% lowers the toughness and weld crack resistance of the base metal and the welded portion, so the upper limit was made 0.3%.
[0028]
Next, Si is an element effective as a deoxidizing element and for securing the strength of the base material. However, if it is less than 0.01%, deoxidation is insufficient, and it is disadvantageous for securing the strength. On the other hand, an excessive content exceeding 2% forms a coarse oxide and causes deterioration of ductility and toughness. Therefore, the range of Si is set to 0.01 to 2%.
[0029]
Mn is an element necessary for ensuring the strength and toughness of the base material, and it is necessary to contain at least 0.1% or more. In order to deteriorate the toughness of the material, the toughness of the welded portion, and the weld cracking property, the upper limit was made 3% within the allowable range of the material.
[0030]
Al is an element effective for deoxidation, refinement of the heated austenite grain size, etc., but in order to exert the effect, it is necessary to contain 0.001% or more. On the other hand, if it exceeds 0.1% and excessively contained, a coarse oxide is formed and the ductility is extremely deteriorated. Therefore, it is necessary to limit it to the range of 0.001% to 0.1%.
[0031]
N is effective in refining austenite grains in combination with Al and Ti, so that it is effective for improving mechanical properties if it is in a very small amount. Further, it is impossible to remove N in steel completely industrially, and reducing it more than necessary is not preferable because it places an excessive load on the manufacturing process. Therefore, the lower limit is set to 0.001% as a range that can be industrially controlled and the load on the manufacturing process is allowable. If excessively contained, solid solution N increases, which may adversely affect ductility and toughness, so the upper limit is made 0.01% as an acceptable range.
[0032]
P is an impurity element and is harmful to various properties of the steel. Therefore, it is preferable to reduce it as much as possible.
[0033]
S is basically an impurity element, and is particularly preferably reduced because it has a significant adverse effect on the ductility, toughness and fatigue properties of steel. For practical purposes, the upper limit is limited to 0.01% as the amount of adverse effects. However, in the trace amount range, S forms fine sulfides and contributes to improvement of the weld heat affected zone (HAZ) toughness. Therefore, when considering HAZ toughness, S is added in a range of 0.0005 to 0.005%. It is preferable.
[0034]
The above is the reason for limiting the basic components of the thick steel plate of the present invention. In the present invention, for the adjustment of strength and toughness, Ni, Cu, Cr, Mo, W, Ti, V, One or more of Nb, Zr, Ta and B can be contained.
[0035]
Ni is a very effective element that can simultaneously improve the strength and toughness of the base material, but it needs to be added in an amount of 0.01% or more in order to exert its effect. As the amount of Ni increases, the strength and toughness of the base material are improved. However, when the addition exceeds 6%, the effect is saturated, but there is a concern that the HAZ toughness and weldability may be deteriorated. Since it is an expensive element, the upper limit of Ni is set to 6% in the present invention in consideration of economy.
[0036]
Cu is an element having almost the same effect as Ni. However, in order to exert the effect, addition of 0.01% or more is necessary, and if it exceeds 1.5%, hot workability and HAZ toughness are required. In the present invention, the content is limited to 0.01 to 1.5%.
[0037]
Cr is an element effective for improving the strength by solid solution strengthening and precipitation strengthening, and 0.01% or more is necessary to produce the effect. However, if Cr is added excessively, the quenching hardness increases and coarse precipitates. Since the base material and the toughness of the HAZ are adversely affected through the formation of the metal, etc., the upper limit is limited to 2% as an acceptable range.
[0038]
Like Cr, Mo and W are effective elements for increasing the strength by solid solution strengthening and precipitation strengthening. However, each of Mo and W is 0 as a range in which the effect can be exerted and other characteristics are not adversely affected. Limited to 0.01-2%.
[0039]
Since Ti forms stable TiN in austenite and contributes not only to the base material but also to the refinement of the heated austenite grain size of HAZ, it is an element effective for improving toughness in addition to improving strength. However, in order to exert its effect, it is necessary to contain 0.003% or more, but when it is contained excessively exceeding 0.1%, coarse TiN is formed and the toughness is deteriorated conversely. In the present invention, the content is limited to a range of 0.003 to 0.1%.
[0040]
V is an element effective for improving the strength of the base material by precipitation strengthening, but it needs to be 0.005% or more in order to exert the effect. As the addition amount increases, the strengthening amount also increases, but with this, the base metal toughness and HAZ toughness deteriorate, and the precipitates become coarser and the strengthening effect tends to be saturated. The upper limit is 0.5% in a range where the toughness deterioration is small.
[0041]
Nb is an element effective for increasing the strength in a small amount by precipitation strengthening and transformation strengthening, and also has a large effect on the processing and recrystallization behavior of austenite, and is therefore effective for improving the toughness of the base metal. Furthermore, it is effective for improving the fatigue characteristics of HAZ. In order to exert the effect, 0.003% or more is necessary. However, if it is added excessively exceeding 0.2%, the toughness is extremely deteriorated. Therefore, in the present invention, it is limited to the range of 0.003 to 0.2%.
[0042]
Zr is also an element effective for improving the strength mainly by precipitation strengthening, but 0.003% or more is necessary to exert the effect. On the other hand, if it is added excessively exceeding 0.1%, coarse precipitates are formed and the toughness is adversely affected, so the upper limit is made 0.1%.
[0043]
Ta has the same effect as Nb, and contributes to the improvement of strength and toughness by adding an appropriate amount. However, if it is less than 0.005%, the effect is not clearly produced, and if it exceeds 0.2%, excessive addition is not possible. Since toughness deterioration due to coarse precipitates becomes significant, the range is made 0.005 to 0.2%.
[0044]
B is an element that enhances hardenability in a very small amount and is effective for increasing the strength. Since B is segregated at the austenite grain boundary in a solid solution state to enhance the hardenability, it is effective even with a very small amount. This is not preferable because the effect of improving the property is insufficient or the effect tends to vary. On the other hand, if added over 0.005%, coarse precipitates are often formed at the time of steel slab production or at the reheating stage, so the effect of improving hardenability becomes insufficient, There is also an increased risk of toughness degradation due to precipitates. Therefore, in the present invention, the range of B is set to 0.0002 to 0.005%.
[0045]
Furthermore, in this invention, 1 type, or 2 or more types of Mg, Ca, and REM can be contained as needed for the improvement of ductility and the improvement of joint toughness.
[0046]
Mg, Ca, and REM are all effective in improving ductility by suppressing the extension of sulfide during hot rolling. It effectively works to improve joint toughness by refining oxides. The lower limit content for exhibiting the effect is 0.0005% for Mg, 0.0005% for Ca, and 0.005% for REM. On the other hand, excessive content causes coarsening of sulfides and oxides, leading to deterioration of ductility, toughness, and fatigue properties. Therefore, the upper limit is 0.01% for Mg and Ca, and 0.1% for REM, respectively. And
[0047]
The above is the reason for limiting the microstructure and chemical composition, which are the basic requirements of the present invention. In addition, in the present invention, an appropriate manufacturing method for satisfying the organizational requirements of the present invention is also presented. However, for the microstructure of the present invention, the effect is exhibited regardless of the means for achieving it, and the method for producing a thick steel sheet having excellent fatigue strength according to claims 1 to 3 of the present invention is described in the claims. It is not limited to the methods shown in 4 and 5.
[0048]
In the production method presented in the steel of the present invention, the steel slab having the chemical composition of the present invention is converted into AC._ThreeReheated to transformation temperature to 1250 ° C, start temperature is 850 ° C or less, end temperature is Ar_ThreeAfter rolling at a transformation point or higher and including a rolling with a cumulative reduction ratio of 30% or more and a total rolling ratio of 5 or more, the steel is cooled to 500 ° C. or less at a cooling rate of 5 ° C./s or less, and then (AC₁Transformation point + 30 ° C) to (AC_ThreeIt is characterized by being reheated to a transformation point of −50 ° C., cooled to 400 ° C. or lower at 5 to 100 ° C./s, and tempered at 250 to 600 ° C. as necessary. Hereinafter, various reasons for limiting the manufacturing method will be described in detail.
[0049]
First, prior to hot rolling, the steel slab is AC_ThreeReheat to transformation point to 1250 ° C. Reheating temperature is AC_ThreeIf it is less than the transformation point, uniform austenitization is not achieved, and the untransformed region remains non-uniformly, which hinders control of the final structure. Moreover, when a precipitate forming element is contained, the solution of the element is not sufficient, and there is a possibility that it is not effectively used for increasing the strength or improving the toughness. On the other hand, when the heating temperature exceeds 1250 ° C., the heated austenite grain size may be excessively coarsened.In this case, the coarse austenite structure is not eliminated even by hot rolling, and in the final structure, the requirements of the present invention, in particular, There is a high possibility that (1) the average ferrite grain size is 20 μm or less and (4) the average aspect ratio of the hard second phase is 10 or more. For the above reasons, in the present invention, the reheating temperature of the steel slab is set to AC._ThreeThe transformation point is limited to 1250 ° C.
[0050]
AC steel billet_ThreeAfter reheating to the transformation point to 1250 ° C., hot rolling is performed. The requirements for hot rolling in the present invention are that the total reduction ratio (steel piece thickness / steel plate thickness) is 5 or more and the start temperature is 850 ° C. or less. The end temperature is Ar_ThreeIt includes rolling with a cumulative rolling reduction of 30% or more above the transformation point. The total reduction ratio of 5 or more is a requirement for controlling the structure of the hard second phase, and the start temperature is 850 ° C. or less and the end temperature is Ar._ThreeThe inclusion of rolling at a transformation point or higher and a cumulative rolling reduction of 30% or more is a necessary condition for controlling the structure of the hard second phase and controlling the ferrite grain size to 20 μm or less. In addition, the steel slab as referred to in the present invention includes an ingot as cast and ingot, a steel slab made by ingot casting into a steel slab, and a steel slab produced by continuous casting.
[0051]
The hard second phase is finally formed in the heat treatment stage after hot rolling, but the region that is austenitized in the heating stage of the heat treatment and transformed into the hard second phase in the subsequent cooling stage is almost the same. Corresponds to the second phase region where C was concentrated during rolling. Therefore, the form of the hard second phase is almost determined not only by the various second phases formed in the rolling stage before the heat treatment, that is, not only by bainite and martensite but also by the form of the second phase including pearlite. The requirements for rolling in the present invention are requirements for ensuring that the aspect ratio of the second phase before heat treatment is sufficiently large and that the aspect ratio of the hard second phase after heat treatment is 10 or more.
[0052]
That is, the second phase formed during hot rolling is formed along the microsegregation with an increased alloy composition if the cooling rate is such that ferrite transformation precedes, and is randomly distributed at the billet stage by rolling. The micro-segregated portion that has been stretched in a band shape parallel to the rolling direction. Moreover, if the state of microsegregation is the same, the band-like structure is more emphasized by promoting the ferrite transformation and concentrating C into untransformed austenite.
[0053]
In order to extend the micro-segregation part in the rolling direction, the total reduction ratio may be specified, the start temperature is 850 ° C. or less, and the end temperature is Ar._ThreeUnder the condition of including the rolling at the transformation point or higher and the cumulative rolling reduction of 30% or higher, and cooling at a cooling rate of 5 ° C./s or lower to 500 ° C. or lower after rolling, if the total rolling ratio is 5 or higher, the final The average aspect ratio of the hard second phase of the structure can be reliably set to 10 or more.
[0054]
Start temperature is 850 ° C or less, end temperature is Ar_ThreeRolling at a transformation point or higher and a cumulative reduction ratio of 30% or more promotes ferrite transformation, so that the average ferrite grain size of the final structure is 20 μm or less, and the band structure at the hot rolling stage is more emphasized. By refining recrystallized austenite under the conditions necessary to achieve this, or introducing processing strain in an unrecrystallized state, the ferrite is refined and ferrite transformation is promoted. The starting temperature is set to 850 ° C. or lower because when the starting temperature is higher than 850 ° C., the recrystallized austenite does not become fine grain or does not become non-recrystallized austenite depending on the chemical composition. This is because the effect of promoting the effect is small. The end temperature is Ar_ThreeThe end temperature is set to Ar or higher than the transformation point._ThreeIf it is below the transformation point, it becomes rolled in the austenite / ferrite two-phase region, resulting in a structure containing processed ferrite, which is not preferable for toughness, and austenite is nucleated randomly from the processed ferrite during the final heat treatment. This is because the formation of band-like tissue is inhibited.
[0055]
Start temperature is 850 ° C or less, end temperature is Ar_ThreeThe cumulative rolling reduction in rolling above the transformation point is 30% or more. If the cumulative rolling reduction is less than 30%, austenite is not sufficiently refined by recrystallization, and in the case of non-recrystallized zone rolling, the amount of work strain is insufficient. Is not sufficient. Further, the ferrite transformation is not sufficiently promoted, and as a result, the formation of a band structure having a large aspect ratio becomes insufficient, and the average aspect ratio of the final hard second phase may not be sufficiently increased.
[0056]
In hot rolling, the start temperature is 850 ° C. or less and the end temperature is Ar._ThreeIf the rolling point is equal to or higher than the transformation point and the cumulative reduction ratio is 30% or more, it is separately over 850 ° C. in order to adjust the sheet thickness and to secure the total reduction ratio, which is a requirement of the present invention. Including hot rolling at is acceptable. If the requirements of the present invention are satisfied, even if rolling above 850 ° C. is included, the object of the present invention is to adversely affect the fatigue strength and the strength and toughness that are the basic properties of structural materials. Absent. On the other hand, the start temperature is 850 ° C. or lower, and the end temperature is Ar_ThreeEven when the rolling point is higher than the transformation point and the cumulative rolling reduction is 30% or more, Ar_ThreeFurther rolling below the transformation point may have an adverse effect, as is apparent from the reasons described above. However, Ar_ThreeRolling below the transformation point is allowed as long as the cumulative rolling reduction is less than 30%, since there is almost no adverse effect on the formation of structural requirements and the material of the present invention.
[0057]
In the production method of the present invention, after the hot rolling, it is necessary to further cool to 500 ° C. or less at a cooling rate of 5 ° C./s or less. The cooling condition is also a condition for clearly forming a band-like structure in the structure in the rolling stage before the final heat treatment. That is, when the cooling rate increases beyond 5 ° C./s, the start temperature is 850 ° C. or less and the end temperature is Ar._ThreeEven if the formation of ferrite is promoted in advance by rolling at a transformation point or higher and a cumulative reduction ratio of 30% or more, depending on the chemical composition, the formation of ferrite is not sufficient, resulting in the formation of a band structure with a large aspect ratio. It is because is inhibited. However, since the transformation of the ferrite and the formation of the band structure are almost completed when the temperature is 500 ° C. or lower, the cooling rate may be controlled to 500 ° C. or lower. Cooling at 500 ° C. or lower may be accelerated cooling, air cooling, or slow cooling.
[0058]
In the production method of the present invention, the steel slab having the chemical composition of the present invention is converted into AC._ThreeReheated to transformation temperature to 1250 ° C, start temperature is 850 ° C or less, end temperature is Ar_ThreeAfter rolling at a transformation point or higher and a rolling reduction of 30% or more in cumulative reduction, the total rolling ratio is 5 or more, and after cooling at a cooling rate of 5 ° C./s or less to 500 ° C. or less, Heating temperature is (AC₁Transformation point + 30 ° C) to (AC_ThreeAt a transformation point of −50 ° C.), heat treatment is performed at a cooling rate of 5 to 100 ° C./s up to 400 ° C. The heat treatment is composed of bainite, martensite, or a mixed structure of both in the final steel sheet structure, and forms 10 to 70% of a hard second phase having an average Vickers hardness of 230 or more and an average aspect ratio of 10 or more. It is a process for doing.
[0059]
By hot rolling prior to heat treatment, a structure in which the average ferrite grain size is 20 μm or less and the second phase becomes a band-like structure having a large aspect ratio is formed, but at this stage, the second phase does not necessarily have fatigue strength. It is not a hard second phase that is effective for improvement and satisfies the present invention. Therefore, the second phase is changed to the one having the necessary characteristics and form while the ferrite is kept fine by heat treatment.
[0060]
In order to make only the second phase a hard second phase using transformation without inducing the excessive coarsening of the fine ferrite formed in the rolling stage, the ferrite is reheated to the two-phase region, By forming a two-phase structure of tote and austenite and cooling at an appropriate cooling rate after reheating, the austenite phase is transformed into a hard second phase composed of bainite, martensite, or a mixed structure of both. At that time, the reheating temperature is (AC₁Transformation point + 30 ° C) to (AC_ThreeTransformation point −50 ° C.). The reheating temperature is (AC₁If the temperature is less than the transformation point + 30 ° C., the ratio of reverse transformation to the austenite phase in the heating stage becomes too small, and it may be impossible to ensure 10% or more of the hard second phase. On the other hand, the reheating temperature is (AC_ThreeIf the transformation point is over 50 ° C), the ratio of a) reverse transformation to the austenite phase in the heating stage becomes excessive, so the ratio of the hard second phase exceeds 70%. The amount of C enrichment in the phase is insufficient, and the hardness of the hard second phase becomes too small. C) The band-like structure of the hard second phase collapses, and the aspect ratio of the hard second phase is 10 or more. D) Various adverse effects such as ferrite grain growth and an average grain size exceeding 20 μm are caused. For the above reasons, in the present invention, the reheating temperature in the final heat treatment is set to (AC₁Transformation point + 30 ° C) to (AC_ThreeTransformation point-50 ° C). (AC₁Transformation point + 30 ° C) to (AC_ThreeThe holding time at the transformation point −50 ° C. does not need to be specified in the industrial range, but holding for an extremely long time is not particularly advantageous for improving the characteristics, and the productivity is lowered. Since there is a concern about the coarsening of the ferrite grain size, it is recommended that the holding time be within about 10 hours.
[0061]
(AC₁Transformation point + 30 ° C) to (AC_ThreeTransformation point −50 ° C.), and the whole or part of the region austenitized in the heating stage, which is a requirement of the present invention, has an average Vickers hardness of 230 or more and an average aspect ratio of In order to transform into a hard second phase composed of 10 or more bainite or martensite or a mixed structure of both, cooling is performed at 5 to 100 ° C./s to 400 ° C. or less. That is, in the chemical composition range of the present invention, a hard second phase having an average Vickers hardness of 230 or more cannot be formed at a cooling rate of less than 5 ° C./s. If the cooling rate is 5 ° C./s or higher, it is preferable that the cooling rate is larger. Therefore, in the present invention, the cooling rate is limited to 5 to 100 ° C./s. The cooling rate must be controlled until the transformation from austenite to bainite or martensite is almost completed. Therefore, in this invention, it shall cool at 5-100 degrees C / s to 400 degrees C or less. For cooling below 400 ° C, the influence on the form of the hard second phase and the hardness is very small, and it is not necessary to limit it in particular. However, if it is extremely slowly cooled around 400 ° C, the toughness due to the formation of coarse precipitates Since there is a risk of promoting deterioration and embrittlement of grain boundaries, it is preferable to avoid extreme slow cooling at an average cooling rate at 400 to 200 ° C. of 1 ° C./min or less.
[0062]
The above is the reason for limiting the basic requirements regarding the manufacturing method in the present invention. However, tempering can be further performed after the two-phase region heat treatment for the purpose of adjusting the strength and reducing the residual stress. However, the tempering temperature is limited to 250 to 600 ° C. This is because if the tempering temperature is less than 250 ° C., the tempering effect is not clear, while if it exceeds 600 ° C., the hard second phase is greatly softened and the Vickers hardness is likely to be less than 230. The tempering holding time and cooling conditions are not particularly required as long as the enjoyment of the tempering effect and the securing of the hardness of the hard second phase are compatible, but when tempering at 500 ° C. or higher, Since there is a great concern about the softening of the second phase, it is preferable to keep the holding time within a short time of 10 hours.
[0063]
Next, the effects of the present invention will be described more specifically with reference to examples.
[0064]
【Example】
Table 1 shows the chemical composition of the test steel used in the examples. Each test steel is made into a steel slab by ingot rolling, by ingot rolling, or by continuous casting. In Table 1, steel slab numbers 1 to 10 satisfy the chemical composition range of the present invention, and steel slab numbers 11 to 15 do not satisfy the chemical composition range of the present invention. Table 1 also shows the heating transformation point (AC₁, AC_ThreeThis indicates that the rate of temperature rise is 5 ° C./min. The measured values at the time are substantially in agreement with the transformation points shown in Table 3 under the actual temperature rise conditions in the heat treatment of the steel sheet.
[0065]
[Table 1]

[0066]
Steel strips having the chemical composition shown in Table 1 were manufactured into steel plates having a thickness of 25 mm or 50 mm by hot rolling under the conditions shown in Table 2, heat treatment under the conditions shown in Table 3, and partly tempering, Tensile properties, 2 mm V notch Charpy impact properties, and fatigue properties of welded joints were investigated. Tensile test pieces and Charpy impact test pieces were sampled perpendicularly to the rolling direction (C direction) from the center of the plate thickness. Tensile properties were measured at room temperature, and Charpy impact properties were evaluated at 50% fracture surface transition temperature (vTrs). The fatigue test was performed on the turn welded joint shown in FIG. 1 in order to evaluate the fatigue characteristics when a fatigue crack occurs from the weld toe of the structure and propagates through the base metal. The test piece is a test plate with a size from steel plate to steel plate longitudinal direction length: 300 mm, width direction length: 80 mm, plate thickness: 25 mm (total thickness for 25 mm thick material, sampled from the surface for 50 mm thick material). A rib plate having a width of 10 mm, a length of 30 mm, and a height of 30 mm was taken around the center of the test plate by carbon dioxide welding (Co2 welding) and welded by welding. In this case, carbon dioxide gas welding uses a 1.4 mm diameter welding wire having a chemical composition of C: 0.06 mass%, Si: 0.5 mass%, Mn: 1.4 mass%, current: 270A, Voltage: 30 V, welding speed: 20 cm / min. I went there. In the fatigue test, the span of the load fulcrum P was set to the lower span: 70 mm, the upper span: 220 mm, the maximum load (Pmax): 5500 kgf, and the stress ratio (R): 0.1 repeated stress load, and the fatigue life was measured. did.
[0067]
[Table 2]

[0068]
[Table 3]

[0069]
Notes in Tables 2 and 3 are as follows.
Note 1) “850 ° C. or lower, end temperature Ar essential for the present invention₃Rolling prior to “rolling with a cumulative reduction of 30% or more”
Note 2) “850 ° C. or lower, end temperature Ar essential for the present invention₃As described above, rolling under a cumulative reduction ratio of 30% or more and rolling under conditions corresponding to the main rolling in the comparative example
Note 3) Measured value by air cooling after rolling
Note 4) Means are air cooling (AC) or accelerated cooling (AcC). In the case of air cooling, cool to room temperature. The cooling rate is the average from the end of rolling to 500 ° C.
Note 5) Average from heating temperature to cooling stop temperature.
Note 6) Tempering retention is in the range of 30 minutes to 2 hours. All cooling is air cooling.
[0070]
Table 4 shows the structure of the steel sheet (ferrite particle size, hard second phase type, ratio, Vickers hardness, aspect ratio) and mechanical properties. In addition, the quantification of the structure was obtained by photographing an optical microscope structure of 1 mm below the surface, 1/4 of the plate thickness, and the center of the plate thickness in the plate thickness section parallel to the longitudinal direction of the steel plate, and the average ferrite particle size at each position, The ratio of the hard second phase and the aspect ratio were determined, and the average value at the three observation positions was shown. The hardness of the hard second phase is 1 mm below the surface, 1/4 of the plate thickness, the center of the plate thickness, and 10 or more micro Vickers hardnesses with a load of 5 to 10 g are measured at each position, and the average value at each position Were further averaged.
[0071]
[Table 4]

[0072]
The steel plate numbers A1 to A13 in Tables 2 to 4 are steel plates that satisfy all the requirements regarding the chemical composition and structure of the present invention, all of which are necessary strength and toughness for structural steel (2 mm V notch Charpy impact properties). It is clear that it also has good joint fatigue properties.
[0073]
On the other hand, steel plate numbers B1 to B11 are comparative steel plates that do not satisfy any of the requirements of the present invention, and it is clear that joint fatigue properties and toughness are inferior to the steel plates of the present invention. .
[0074]
Steel plate numbers B1 to B5 have good characteristics even though they cannot satisfy the structural requirements of the present invention because the chemical composition does not satisfy the present invention, or satisfy the structural requirements of the present invention. This is an example that could not be achieved.
[0075]
That is, the steel plate number B1 has an excessively small amount of C, so the hard second phase has low hardenability, and the entire surface has a mixed structure of pearlite and bainite. The hardness of the hard second phase is lower than the lower limit of the present invention. Is also low. Therefore, compared with the steel plate of the present invention, joint fatigue characteristics are greatly inferior.
[0076]
Steel plate number B2 is a comparative example in which the amount of C is excessive, and the toughness is extremely inferior because the hardness of the hard second phase is excessive and very brittle. Although the joint fatigue properties are relatively good, the joint fatigue properties are slightly inferior to those of the present invention due to the brittle fracture of the hard phase in fatigue tests.
[0077]
Steel plates Nos. B3 and B4 have excessive amounts of Mn and P, respectively. Further, due to the deterioration of toughness, in the joint fatigue test, brittle fracture in the ferrite phase and the hard phase partially occurs, so the joint fatigue characteristics are inferior to those of the present invention.
[0078]
Steel plate number B5 is also inferior in toughness and joint fatigue properties compared to the present invention because the amount of S is excessive. In particular, the deterioration of joint fatigue characteristics is large.
[0079]
Steel plate numbers B6 to B11 are examples in which the chemical composition satisfies the present invention, but the structural requirements do not satisfy the present invention, so the joint fatigue characteristics are inferior or the toughness is not sufficient as structural steel. is there.
[0080]
That is, the steel plate number B6 has a final heating temperature of AC_ThreeSince it exceeds the upper limit of the transformation point and the production method of the present invention, the ferrite phase does not remain, the structure after accelerated cooling becomes a mixed structure of bainite and martensite on the entire surface, and the two phases of ferrite and hard second phase. Since it is not a phase structure, no improvement in joint fatigue characteristics is observed.
[0081]
On the other hand, the steel plate number B7 has a heating temperature of the final heat treatment lower than the lower limit of the production method of the present invention, so that a certain amount of austenite necessary for forming the hard phase is not formed in the heating stage, and as a result, In addition, a hard second phase composed of bainite and martensite essential in the present invention and having a hardness of 230 or more is not formed, and the joint fatigue characteristics are extremely low as compared with the present invention.
[0082]
Steel plate number B8 has a small total reduction ratio, so that the aspect ratio of the hard second phase is too small to satisfy the present invention, and the joint fatigue characteristics are inferior to the present invention.
[0083]
Steel plate No. B9 has a low cooling rate in the final heat treatment, so that the hardness of the hard second phase is too small to function effectively in suppressing crack propagation, and therefore the joint fatigue properties are significantly inferior.
[0084]
Steel plate number B10 is not rolled in hot rolling starting from 850 ° C. or less, and the ferrite grain boundary is excessively out of the scope of the present invention. Therefore, the aspect ratio of the hard second phase is within the scope of the present invention compared to the steel of the present invention rolled at the same total reduction ratio. Is sufficiently superior to the comparative example, but is slightly inferior to the steel of the present invention with other structural requirements that are similar.
[0085]
Steel plate number B11 was subjected to accelerated cooling with a high cooling rate after hot rolling, so that the second phase was hardly layered at the end of rolling, so the aspect ratio of the hard second phase after the final heat treatment was The joint fatigue characteristics are not improved.
[0086]
From the above examples, it is apparent that according to the present invention, it is possible to obtain excellent joint fatigue characteristics while ensuring sufficiently high toughness as structural steel.
[0087]
【The invention's effect】
The present invention is specially used for welded structural members that require fatigue strength, and has specially improved joint fatigue characteristics, which have been conventionally difficult to improve in welded parts, for thick steel sheets with a tensile strength of 400 MPa or higher. The industrial utility is extremely large in that it can be manufactured without resorting to alloying elements or complicated manufacturing processes.
[Brief description of the drawings]
FIG. 1 is a conceptual diagram of a turn welding four-point bending test method for evaluating weld joint fatigue characteristics.
[Explanation of symbols]
P: Load fulcrum

Claims (4)

% By mass
C: 0.04-0.3%
Si: 0.01-2%
Mn: 0.1 to 3%
Al: 0.001 to 0.1%,
N: 0.001 to 0.01%
Containing
P: 0.02% or less,
S: 0.01% or less
Steel, the balance of which is composed of iron and inevitable impurities, is reheated to an AC ₃ transformation point to 1250 ° C., the start temperature is 850 ° C. or less, the end temperature is the Ar ₃ transformation point or more, and the cumulative reduction rate Includes rolling at 30% or more, hot rolling at a total reduction ratio of 5 or more, cooling to 500 ° C. or less at a cooling rate of 5 ° C./s or less, and then (AC ₁ transformation point + 30 ° C.) to ( AC ₃ transformation point-50 ° C.), cooled to 400 ° C. or less at 5-100 ° C./s , having a structure composed of ferrite and a hard second phase, and a plate thickness parallel to the longitudinal direction of the steel plate The ferrite and the hard second phase in the cross-sectional structure satisfy all the following conditions (1) to (4), and the structure of the hard second phase is either bainite, martensite, or a mixed structure of both. A method for producing thick steel plates with excellent fatigue strength
(1) Average ferrite particle size: 20 μm or less
(2) Ratio of hard second phase: 10 to 70%
(3) Average Vickers hardness of hard second phase: 230 or more
(4) Average aspect ratio of hard second phase (average length in the longitudinal direction of steel sheet / average length in the thickness direction): 10 or more Furthermore, the steel slab is mass%,
Ni: 0.01-6%,
Cu: 0.01 to 1.5%,
Cr: 0.01-2%
Mo: 0.01-2%
W: 0.01-2%
Ti: 0.003 to 0.1%,
V: 0.005-0.5%
Nb: 0.003 to 0.2%,
Zr: 0.003 to 0.1%,
Ta: 0.005 to 0.2%,
B: 0.0002 to 0.005%
The manufacturing method of the thick steel plate excellent in the fatigue strength of Claim 1 characterized by including 1 type, or 2 or more types of these. Furthermore, the billet is mass%
Mg: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
REM: 0.005-0.1%
The manufacturing method of the thick steel plate excellent in the fatigue strength of Claim 1 or 2 characterized by including 1 type, or 2 or more types of these . Furthermore, tempering at 250-600 degreeC, The manufacturing method of the thick steel plate excellent in the fatigue strength in any one of Claims 1-3 characterized by the above-mentioned.

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