JP4926406B2 - Steel sheet with excellent fatigue crack propagation characteristics - Google Patents
Steel sheet with excellent fatigue crack propagation characteristics Download PDFInfo
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- JP4926406B2 JP4926406B2 JP2005016036A JP2005016036A JP4926406B2 JP 4926406 B2 JP4926406 B2 JP 4926406B2 JP 2005016036 A JP2005016036 A JP 2005016036A JP 2005016036 A JP2005016036 A JP 2005016036A JP 4926406 B2 JP4926406 B2 JP 4926406B2
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- 229910000831 Steel Inorganic materials 0.000 title claims description 61
- 239000010959 steel Substances 0.000 title claims description 61
- 229910000734 martensite Inorganic materials 0.000 claims description 59
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Classifications
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
- B21B—ROLLING OF METAL
- B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
- B21B3/02—Rolling special iron alloys, e.g. stainless steel
Description
本発明は、疲労特性が必要とされる建築、造船、橋梁、建設機械、海洋構造物などの溶接構造部材に使用される疲労き裂伝播特性に優れた鋼板に関するものである。 The present invention is, architecture fatigue properties are required, shipbuilding, bridges, construction machinery, it relates to a steel plate having excellent fatigue crack propagation properties for use in welded structural members such as marine structures.
一般に、建築、造船、橋梁、建設機械、海洋構造物などの溶接構造物には、アーク溶接、プラズマ溶接をはじめ、レーザ溶接や電子ビーム溶接など、多種多様な溶接方法を用いた溶接継手が適用されている。 Generally, welded joints using a wide variety of welding methods such as arc welding, plasma welding, laser welding, and electron beam welding are applied to welded structures such as buildings, shipbuilding, bridges, construction machinery, and offshore structures. Has been.
これらの溶接継手には、風や波、機械振動などによる繰り返し荷重がかかるため、疲労強度の向上が極めて重要であり、一般的には疲労強度の向上手法として溶接後処理である、（１）グラインディング、（２）ＴＩＧドレッシング、（３）ショットピーニング、（４）ハンマーピーニングが用いられるが、以下のような問題点があった。
ここで、グラインディング、ＴＩＧドレッシングは、溶接ビードの形状をよくするものであるが、いずれも著しく作業効率が悪かった。
ショットピーニング、ハンマーピーニングは、疲労強度向上効果はあるが、ショットピーニングは巨大な機械が必要であるうえ、種々のユーティリティが必要となる。
Since these welded joints are subjected to repeated loads due to wind, waves, mechanical vibrations, etc., it is extremely important to improve fatigue strength. Generally, this is a post-welding process as a method for improving fatigue strength. (1) Grinding, (2) TIG dressing, (3) shot peening, and (4) hammer peening are used, but have the following problems.
Here, the grinding and the TIG dressing improve the shape of the weld bead, but the working efficiency is remarkably poor.
Shot peening and hammer peening have an effect of improving fatigue strength, but shot peening requires a huge machine and various utilities.
またハンマーピーニングは反動が大きく、処理結果が安定せず、時にはかえってプレス成形性や疲労強度を低下させてしまうことがある。またハンマーピーニングは、あまりに大きな塑性変形を与えるために、薄い板に対しては使いにくいという欠点もあった。
さらに、グラインディングやハンマーピーニングは、数Ｈｚの低周波の機械加工を継手部に施すため加工表面の凹凸が激しく、その山部に応力が集中し、継手部に繰り返し荷重がかかると、この応力集中部からき裂が生じるため継手全体の疲労強度が低下するという問題点があった。
In addition, hammer peening has a large reaction, the processing result is not stable, and sometimes press formability and fatigue strength are lowered. Hammer peening also has a drawback that it is difficult to use for thin plates because it gives too much plastic deformation.
Furthermore, since grinding and hammer peening are performed on the joints with low frequency machining of several Hz, the unevenness of the processed surface is severe, stress concentrates on the peaks, and repeated stress is applied to the joints. There is a problem that the fatigue strength of the entire joint is lowered because cracks are generated from the concentrated portion.
また溶接部には、一般に溶接による入熱によって残留応力が導入される。その残留応力が溶接部で疲労強度を低下させる一つの大きな要因となっている。
そこで、疲労強度を向上させる別の手段として、溶接継手部に圧縮残留応力を発生させるか、あるいは溶接継手部に発生する引張残留応力を低減して疲労強度を高める方法が知られている。
例えば、溶接止端部近傍にショットピーニング処理を行うことで圧縮残留応力を付与できる。ここにショットピーニング処理は、疲労き裂発生の起点となる部位に、１ｍｍ弱の鋼球を多数打ち付け圧縮残留応力を付与する手法である。
さらに、溶接金属の加熱再溶融により溶接止端部形状の改善あるいは引張残留応力の軽減が可能であることも知られている。
しかし、このショットピーニング処理は鋼球を必要とし、この鋼球の後処理あるいはコストが問題となる場合がある。さらに疲労強度の向上代がばらつくという問題点がある。
Further, residual stress is generally introduced into the weld by heat input by welding. The residual stress is one major factor that reduces the fatigue strength at the weld.
Therefore, as another means for improving the fatigue strength, a method is known in which a compressive residual stress is generated in the welded joint portion or a tensile residual stress generated in the welded joint portion is reduced to increase the fatigue strength.
For example, compressive residual stress can be applied by performing shot peening near the weld toe. Here, the shot peening treatment is a technique in which a compression residual stress is applied by hitting a large number of steel balls of less than 1 mm at the site where fatigue cracks start.
It is also known that the weld toe shape can be improved or the tensile residual stress can be reduced by remelting the weld metal by heating.
However, this shot peening process requires a steel ball, and the post-treatment or cost of this steel ball may be a problem. Furthermore, there is a problem that the amount of improvement in fatigue strength varies.
以上のように、溶接後処理による疲労強度の向上技術を溶接継手に採用することは困難であり、たとえ採用できても疲労強度向上代が低いレベルに留まっていた。そこで溶接後処理が不要であり、溶接ままで溶接継手の疲労強度向上が達成できるような技術が切望されている。 As described above, it is difficult to employ a technique for improving fatigue strength by post-welding treatment in a welded joint, and even if it can be adopted, the allowance for improving fatigue strength remains at a low level. Therefore, there is a need for a technique that does not require post-weld processing and that can achieve improved fatigue strength of the welded joint as it is.
このような観点から、溶接ままで溶接継手の疲労強度を向上させるために、疲労き裂の伝播を抑制する鋼板がいくつか提案されている。
例えば非特許文献１には、一般造船用鋼材の昇温過程でフェライトを加工することにより表層に超細粒組織を形成したいわゆるＳＵＦ鋼が、疲労き裂の伝播速度を低下させる効果を有することが開示されている。しかし、フェライトの細粒化のみでは伝播速度を著しく低下させるのは困難であり、さらに表層に形成された超細粒組織は溶接熱影響により大部分消失してしまうので、溶接継手の疲労強度向上は十分達成できていない。
From this point of view, several steel sheets that suppress the propagation of fatigue cracks have been proposed in order to improve the fatigue strength of the welded joint as it is.
For example, Non-Patent Document 1 discloses that so-called SUF steel in which a superfine grain structure is formed on the surface layer by processing ferrite in the temperature rising process of general shipbuilding steel has an effect of reducing the propagation speed of fatigue cracks. Is disclosed. However, it is difficult to significantly reduce the propagation speed only by reducing the ferrite grain size, and the superfine grain structure formed on the surface layer is largely lost due to the effect of welding heat, improving the fatigue strength of welded joints. Is not fully achieved.
また特許文献１〜７には、軟質のフェライト母相中に硬質のパーライト、ベイナイト、マルテンサイトを第二相とした混合組織にすると、硬質第二相がき裂進展の障害となり、疲労き裂の伝播速度を低下することができる鋼板が開示されている。しかしこれらの技術には、き裂進展を遅らせるための重要な要素であるマルテンサイトの面積率、アスペクト比（長軸／短軸）、硬さ、およびフェライトの硬さ、および両者の間隔を適切に制御できていないために、全く疲労き裂伝播特性が向上しない場合や向上が不十分な場合や著しく鋼材の靭性が劣化する場合があった。 Further, in Patent Documents 1 to 7, when a mixed structure having hard pearlite, bainite, and martensite as a second phase in a soft ferrite matrix is used, the hard second phase becomes an obstacle to crack propagation, and fatigue cracks are caused. A steel sheet capable of reducing the propagation speed is disclosed. However, for these technologies, the martensite area ratio, aspect ratio (major axis / minor axis), hardness, and ferrite hardness, which are important factors for delaying crack growth, are appropriately adjusted. Therefore, the fatigue crack propagation characteristics are not improved at all, the improvement is insufficient, or the toughness of the steel material is significantly deteriorated.
例えば特許文献１では、マルテンサイト分率が不十分であり、十分な疲労き裂伝播特性の向上が得られない。特許文献２では、マルテンサイト分率が３０％を超えると著しい靭性の低下が起こる上に、フェライトに対する硬質第二相の硬度を３０％以上確保してもフェライトの硬さが１５０以下や硬質第二相の硬さが４００以下では、疲労き裂伝播特性の十分な向上効果は得られない。特許文献３も同様にマルテンサイト分率が３０％を超えるものであり、鋼材の靭性が著しく損なわれてしまう。
特許文献４〜７では、フェライト、第二相の硬さ、分率、およびそれらの間隔が適切に制御されておらず、第二相が硬さの低い４００以下のベイナイトの場合では、分率が多くても靭性劣化は抑制されるが、伝播抑制効果は小さい。また、第二相が硬さの高い４００以上のマルテンサイトの場合では、分率が３０％以上では著しい靭性劣化が生じる。
For example, in Patent Document 1, the martensite fraction is insufficient, and sufficient improvement in fatigue crack propagation characteristics cannot be obtained. In Patent Document 2, when the martensite fraction exceeds 30%, a significant decrease in toughness occurs, and even if the hardness of the hard second phase with respect to ferrite is ensured to be 30% or more, the hardness of the ferrite is 150 or less. When the two-phase hardness is 400 or less, a sufficient improvement effect of fatigue crack propagation characteristics cannot be obtained. Similarly, Patent Document 3 has a martensite fraction exceeding 30%, and the toughness of the steel material is significantly impaired.
In Patent Documents 4 to 7, the ferrite, the hardness of the second phase, the fraction, and the interval between them are not properly controlled, and in the case of bainite having a low hardness of 400 or less, the fraction Even if there are many, toughness deterioration is suppressed, but the propagation suppression effect is small. Further, in the case where the second phase is 400 or more martensite having a high hardness, if the fraction is 30% or more, significant toughness deterioration occurs.
また特許文献８には、フェライトとベイナイトの二相組織とした上で、フェライト相部分の比率、フェライトの硬さ、およびフェライトとベイナイトの相境界の数等を特定範囲に規定することによって、疲労き裂進展速度を低下することができる鋼板が開示されている。しかし、ベイナイトの硬さレベルでは疲労き裂伝播特性の向上効果は不十分であり、フェライトの硬さが１５０以下でも同様に効果は小さい。 Further, in Patent Document 8, after a two-phase structure of ferrite and bainite is set, the ratio of the ferrite phase portion, the hardness of the ferrite, the number of phase boundaries between the ferrite and bainite, and the like are specified in a specific range, thereby fatigue. A steel sheet capable of reducing the crack growth rate is disclosed. However, the improvement effect of fatigue crack propagation characteristics is insufficient at the bainite hardness level, and the effect is similarly small even if the hardness of the ferrite is 150 or less.
また特許文献９〜１１では、上記に挙げた思想とは異なり、硬質相を母相とし、軟質相を第二相とし分散させることで疲労き裂伝播速度を低下させることができる鋼板が開示されている。これらは、き裂進展に必要な塑性変形エネルギーを軟質相で吸収させることで、き裂閉口挙動を促進させ、き裂進展を抑制させることを狙ったものであるが、溶接引張残留応力が存在する溶接継手では、き裂が容易に開口するので、き裂閉口効果のみでは十分な疲労き裂伝播特性の向上効果は得られない。 Further, Patent Documents 9 to 11 disclose a steel sheet capable of reducing the fatigue crack propagation rate by dispersing the hard phase as a parent phase and the soft phase as a second phase, unlike the ideas listed above. ing. These are intended to absorb the plastic deformation energy necessary for crack growth in the soft phase, thereby promoting crack closing behavior and suppressing crack growth, but there is weld tensile residual stress. In such a welded joint, the crack is easily opened, so that a sufficient fatigue crack propagation characteristic improvement effect cannot be obtained only by the crack closing effect.
さらに特許文献１２、１３では、回復あるいは再結晶フェライト分率を確保し、さらに特定の集合組織を発達させることで疲労き裂伝播速度を低下させることができる鋼板が開示されている。これは特定の集合組織によりき裂進展時のき裂先端の塑性変形を抑制しようとするものであるが、第二相組織が規定されてないフェライトの集合組織のみでは十分な疲労き裂伝播特性は得られない上に、き裂先端の塑性変形は極低ΔＫ領域でしか抑制することができないため適用範囲が著しく狭い。
Further,
以上のように、従来技術では、き裂進展を著しく抑制するための適切な組織制御はできておらず、安定的に疲労き裂伝播速度を低下させることができる鋼板、さらに溶接継手の疲労寿命向上に寄与できる鋼板の開発が切望されている。
本発明の課題は、前述のような従来技術の問題点を解決し、建築、造船、橋梁、建設機械、海洋構造物などの溶接構造部材に使用される疲労き裂伝播特性に優れた鋼板を提供することにある。具体的には、応力比０．１の応力拡大係数範囲ΔＫが２０ＭＰａ√ｍのときの疲労き裂伝播速度ｄａ／ｄＮが１０−８ｍ／ｃｙｃｌｅ以下であり、応力比０．１の入熱量が１０〜３０ｋＪ／ｍｉｎの溶接継手軸力疲労試験のときの疲労寿命が従来鋼の２倍以上を満足する鋼板を提供するものである。
An object of the present invention is to solve the problems of conventional techniques as described above, building, shipbuilding, bridges, construction machinery, marine structures Fatigue Crack Propagation excellent steel plate used in the welding structural members, such as Is to provide. Specifically, the fatigue crack propagation rate da / dN when the stress intensity factor range ΔK with a stress ratio of 0.1 is 20 MPa√m is 10 −8 m / cycle or less, and the heat input with the stress ratio of 0.1. There is to provide a steel plate that fatigue life when the welded joint axial
本発明は前述の課題を解決するために鋭意検討の結果なされたものであり、その要旨とするところは以下の通りである。
（１） 質量％で、
Ｃ ：０．０３〜０．２％、 Ｓｉ：０．０１〜１．６％、
Ｍｎ：０．５〜２％、 Ｐ ：０．０２％以下、
Ｓ ：０．０２％以下、 Ａｌ：０．００１〜０．１％、
Ｎ ：０．００１〜０．００８％
を含有し、残部Ｆｅおよび不可避不純物からなり、母材のミクロ組織が、ビッカース硬さが１５０以上のフェライトを母相とし、ビッカース硬さが４００〜９００、面積率が５〜３０％、アスペクト比（長軸／短軸）が３以上の扁平なマルテンサイトを第二相とした層状組織であり、さらにフェライトとマルテンサイトの板厚方向の平均層間隔が３〜５０μｍであり、応力比０．１の応力拡大係数範囲ΔＫが２０ＭＰａ√ｍのときの疲労き裂伝播速度ｄａ／ｄＮが１０−８ｍ／ｃｙｃｌｅ以下であることを特徴とする、疲労き裂伝播特性に優れた鋼板。
（２） さらに質量％で、
Ｃｕ：０．１〜２．５％、 Ｎｉ：０．１〜５％、
Ｃｒ：０．０１〜１．５％、 Ｍｏ：０．０１〜１．５％、
Ｗ ：０．０１〜１．５％、 Ｔｉ：０．００１〜０．０５％、
Ｎｂ：０．００５〜０．２％、 Ｚｒ：０．００５〜０．２％、
Ｖ ：０．００５〜０．２％、 Ｂ ：０．０００２〜０．００５％
の１種または２種以上を含有することを特徴とする、前記（１）に記載の疲労き裂伝播特性に優れた鋼板。
（３） さらに質量％で、
Ｍｇ：０．０００５〜０．０１％、 Ｃａ：０．０００５〜０．０１％、
ＲＥＭ：０．００５〜０．０５％
の１種または２種以上を含有することを特徴とする、前記（１）または（２）に記載の疲労き裂伝播特性に優れた鋼板。
The present invention has been made as a result of intensive studies in order to solve the above-described problems, and the gist thereof is as follows.
(1) In mass%,
C: 0.03-0.2%, Si: 0.01-1.6%,
Mn: 0.5-2%, P: 0.02% or less,
S: 0.02% or less, Al: 0.001-0.1%,
N: 0.001 to 0.008%
Containing the balance Fe and unavoidable impurities, and the microstructure of the base material is a ferrite having a Vickers hardness of 150 or more, the Vickers hardness is 400 to 900, the area ratio is 5 to 30%, and the aspect ratio (Major axis / minor axis) is a layered structure having flat martensite of 3 or more as a second phase, and the average layer spacing in the plate thickness direction of ferrite and martensite is 3 to 50 μm, and the stress ratio is 0. A steel plate having excellent fatigue crack propagation characteristics, wherein a fatigue crack propagation rate da / dN when a stress intensity factor range ΔK of 1 is 20 MPa√m is 10 −8 m / cycle or less.
(2) Furthermore, in mass%,
Cu: 0.1 to 2.5%, Ni: 0.1 to 5%,
Cr: 0.01 to 1.5%, Mo: 0.01 to 1.5%,
W: 0.01-1.5%, Ti: 0.001-0.05%,
Nb: 0.005 to 0.2%, Zr: 0.005 to 0.2%,
V: 0.005-0.2%, B: 0.0002-0.005%
1 or 2 types or more, The steel plate excellent in the fatigue crack propagation characteristic as described in said (1) characterized by the above-mentioned.
(3) Furthermore, in mass%,
Mg: 0.0005 to 0.01%, Ca: 0.0005 to 0.01%,
REM: 0.005 to 0.05%
The steel plate excellent in fatigue crack propagation characteristics as described in (1) or (2) above, comprising one or more of the following.
本発明によれば、建築、造船、橋梁、建設機械、海洋構造物などの溶接構造部材に使用される疲労き裂伝播特性に優れた鋼板を提供することができる。
具体的には、応力比０．１の応力拡大係数範囲ΔＫが２０ＭＰａ√ｍのときの疲労き裂伝播速度ｄａ／ｄＮが１０−８ｍ／ｃｙｃｌｅ以下であり、溶接継手疲労寿命を従来の２倍以上に向上させることができ、溶接鋼構造物の疲労破壊に対する信頼性を向上させるなど、産業上有用な著しい効果を奏する。
According to the present invention, construction, shipbuilding, bridges, construction machinery, it is possible to provide a superior steel plate fatigue crack propagation properties for use in welded structural members such as marine structures.
Specifically, the fatigue crack propagation rate da / dN when the stress intensity factor range ΔK with a stress ratio of 0.1 is 20 MPa√m is 10 −8 m / cycle or less, and the fatigue life of the welded joint is 2 It can be improved more than twice, and there are significant industrially useful effects such as improving the reliability against fatigue fracture of welded steel structures.
一般的には疲労き裂伝播速度は、鋼材の組織や強度に依存しないことが知られている。しかし、本発明者らは鋭意検討を重ねた結果、フェライトを母相とし、マルテンサイトを第二相として層状分散し、さらに硬さ、面積率、アスペクト比（長軸／短軸）、各相の板厚方向の層間隔を適正に制御することによって疲労き裂伝播速度が従来に比べ著しく低下することを知見した。
疲労き裂伝播速度が低下するメカニズムは、鋼板圧延冷却中にマルテンサイト変態したときに生じるマルテンサイト周囲の内部応力の変化によるものであり、き裂進展に対する駆動力を下げる効果がある。この効果により、マルテンサイト直上においてき裂は停滞し、さらにマルテンサイトの内部を容易に進展することはできず、マルテンサイトの界面に沿ってき裂は迂回したり分岐したりする。このようなき裂停滞による遅延、き裂迂回・分岐による伝播距離の増大、さらに、き裂迂回・分岐に伴う著しいき裂閉口挙動の発現が、疲労き裂伝播速度の大幅な低下を可能とした。
In general, it is known that the fatigue crack propagation rate does not depend on the structure and strength of the steel material. However, as a result of intensive investigations, the present inventors have conducted a layer dispersion with ferrite as a parent phase and martensite as a second phase, and further, hardness, area ratio, aspect ratio (major axis / minor axis), each phase It was found that the fatigue crack propagation rate is significantly reduced by controlling the layer spacing in the thickness direction of the steel.
The mechanism by which the fatigue crack propagation rate decreases is due to the change in internal stress around martensite that occurs when martensite transformation occurs during steel sheet rolling cooling, and has the effect of reducing the driving force for crack propagation. Due to this effect, the crack stagnates just above the martensite and cannot easily propagate inside the martensite, and the crack circumvents or branches along the martensite interface. Such a delay due to crack stagnation, an increase in propagation distance due to crack detouring and branching, and a significant crack closing behavior associated with crack detouring and bifurcation enabled a significant decrease in fatigue crack propagation rate. .
ミクロ組織の限定範囲の理由を以下に述べる。
疲労き裂伝播速度に影響する因子の中で最も影響が大きいものは、マルテンサイト面積率であり、５％以上で伝播速度は急激に低下する。これは、マルテンサイト分率が増えることによりき裂進展の障害が増えることに起因する。しかし３０％超では著しく靭性が劣化するので、その範囲を５〜３０％とした。
The reason for the limited range of the microstructure will be described below.
Among the factors affecting the fatigue crack propagation rate, the most influential factor is the martensite area ratio, and the propagation rate rapidly decreases at 5% or more. This is due to the increase in crack propagation failure due to an increase in the martensite fraction. However, if it exceeds 30%, the toughness deteriorates remarkably, so the range was made 5-30%.
内部応力を高めて、き裂進展の駆動力をより効果的に下げるためにはマルテンサイト変態開始温度を低くする必要がある。これは低温でマルテンサイト変態が起こると、変態の拘束となるフェライトが硬いため、その反力により内部応力が増大するからである。
マルテンサイト変態開始温度は、熱間圧延時のオーステナイト中の炭素が濃縮する量が多いほど低下する。そして、炭素が濃縮する量が多いほどマルテンサイトの硬さは大きくなるので、マルテンサイト変態開始温度を４００℃以下にするためマルテンサイトの硬さは４００以上にしなければならない。
マルテンサイト変態開始温度を４００℃以下にしなければならない理由は、４００℃超では変態後の熱収縮により内部応力が緩和され、疲労き裂伝播遅延効果が小さくなるからである。また、マルテンサイトの硬さが９００超では、マルテンサイト分率を５％以上確保することが困難であるとともに、マルテンサイトが起点となり脆性破壊を起こす可能性があるため、マルテンサイトの硬さは４００〜９００Ｈｖとした。
In order to increase internal stress and lower the driving force for crack propagation more effectively, it is necessary to lower the martensitic transformation start temperature. This is because when martensitic transformation occurs at a low temperature, the ferrite acting as a constraint of the transformation is hard, and the internal stress is increased by the reaction force.
The martensitic transformation start temperature decreases as the amount of carbon in the austenite during hot rolling increases. And since the hardness of a martensite becomes so large that there is much quantity which carbon concentrates, in order to make a martensite transformation start temperature into 400 degrees C or less, the hardness of a martensite must be 400 or more.
The reason why the martensitic transformation start temperature must be 400 ° C. or lower is that if it exceeds 400 ° C., the internal stress is relaxed by the thermal contraction after transformation, and the fatigue crack propagation delay effect is reduced. In addition, when the hardness of martensite exceeds 900, it is difficult to secure a martensite fraction of 5% or more, and since martensite may be the starting point and cause brittle fracture, the hardness of martensite is It was set to 400 to 900 Hv.
さらに、上記に述べたようにフェライトの硬いほどマルテンサイト変態時の拘束となり反力が大きくなり内部応力が高まるため、フェライトの硬さを１５０Ｈｖ以上とした。 Furthermore, as described above, the harder the ferrite, the more constrained it is during martensitic transformation, and the reaction force increases and the internal stress increases. Therefore, the hardness of the ferrite is set to 150 Hv or more.
マルテンサイトのアスペクト比が大きいほど、き裂進展の障害となるマルテンサイトに当たる頻度が増える上に、迂回・分岐距離が増大することから疲労き裂伝播速度低下には有効である。アスペクト比が３より小さいと、き裂がマルテンサイトに当たっても迂回・分岐距離が小さいため、き裂伝播特性向上効果は小さい。従ってマルテンサイトのアスペクト比（長軸／短軸）は３以上とした。 The larger the martensite aspect ratio, the more frequently it hits martensite, which is an obstacle to crack growth, and the detour / branch distance increases, which is effective in reducing the fatigue crack propagation rate. When the aspect ratio is less than 3, the detour / branch distance is small even when the crack hits martensite, and the effect of improving crack propagation characteristics is small. Therefore, the martensite aspect ratio (major axis / minor axis) was set to 3 or more.
フェライト相とマルテンサイト相は層状に分散させる必要があり、その層間隔が３μｍより小さいとマルテンサイト変態時に導入される内部応力が有効に働かなくなり、き裂進展を遅延させることは困難となる。また層間隔が５０μｍ超では、き裂がマルテンサイトへ当たる頻度、すなわち、き裂の停滞、迂回・分岐効果が小さくなることから、層間隔の範囲を３〜５０μｍとした。 It is necessary to disperse the ferrite phase and the martensite phase in layers, and if the layer spacing is smaller than 3 μm, the internal stress introduced at the time of martensite transformation does not work effectively, and it becomes difficult to delay the crack growth. In addition, when the layer interval exceeds 50 μm, the frequency with which the crack hits martensite, that is, the stagnation of the crack and the detour / branch effect becomes small, so the range of the layer interval is set to 3 to 50 μm.
次に、各合金元素の範囲を限定した理由を以下に述べる。なお、以下において％は質量％を意味する。
Ｃは、本発明の成分として主たる元素の一つであり、マルテンサイト分率を制御することと、鋼の強度を向上させる有効な成分として含有するもので、０．０３％未満ではマルテンサイト分率を５％以上確保するのが困難であるが、下限は、実施例の表１に示す本発明例３及び１３のＣ量に基づいて０．０５％とした。０．２％超では母材および溶接部の靭性や耐溶接割れ性を低下させるので、０．０５〜０．２％とした。
Next, the reason for limiting the range of each alloy element will be described below. In the following,% means mass%.
C is one of the main elements as a component of the present invention, and is contained as an effective component for controlling the martensite fraction and improving the strength of the steel. When it is less than 0.03%, the martensite component is contained. Although it is difficult to secure a rate of 5% or more , the lower limit was set to 0.05% based on the amount of C in Invention Examples 3 and 13 shown in Table 1 of the Examples . In 0.2 percent since lowering the toughness and resistance to weld cracking resistance of the base metal and weld portion was 0.0 5 to 0.2 percent.
Ｓｉは、強度確保のほか脱酸元素として必須の元素であり、その効果を得るためには０．０１％以上の添加が必要で、１．６％を超える過剰な含有は粗大な酸化物を形成して延性や靭性の低下を招くが、上限は、実施例の表１に示す本発明例１１のＳｉ量に基づいて０．８％とし、その量を０．０１〜０．８％とした。 In addition to ensuring strength, Si is an essential element as a deoxidizing element, and in order to obtain the effect, addition of 0.01% or more is necessary, and excessive content exceeding 1.6% is a coarse oxide. The upper limit is 0.8% based on the Si amount of Invention Example 11 shown in Table 1 of the Examples , and the amount is 0.01 to 0.8 %. did.
Ｍｎは、強度を高めるために必須の元素であるが、０．５％未満では母材強度を確保できない。一方、２％を超える過剰な含有は、粒界脆化等により母材靭性や溶接部の靭性、さらに溶接割れ性などを劣化させるため、その量を０．５〜２％とした。 Mn is an essential element for increasing the strength, but if it is less than 0.5%, the strength of the base material cannot be secured. On the other hand, an excessive content exceeding 2% deteriorates the base metal toughness, the toughness of the welded portion, the weld cracking property, etc. due to grain boundary embrittlement or the like.
Ｐは、鋼の靭性に影響を与える元素であり、０．０２％を超えると母材だけでなくＨＡＺの靭性を著しく阻害するので極力少ないほうが良く、上限を０．０２％とした。 P is an element that affects the toughness of the steel. If it exceeds 0.02%, not only the base metal but also the toughness of HAZ is significantly inhibited.
Ｓは、Ｐと同様に低いほど好ましく、０．０２％を超えるとＭｎＳ析出が顕著となり、母材のＨＡＺ靭性を阻害して板厚方向の延性も低下させるため、上限を０．０２％とした。 S is preferably as low as P, and when it exceeds 0.02%, MnS precipitation becomes prominent, and the HAZ toughness of the base metal is inhibited to reduce the ductility in the thickness direction. did.
Ａｌは、脱酸、オーステナイト粒径の細粒化等に有効な元素であり、効果を発揮するためには０．００１％以上含有する必要がある。一方、０．１％を超えて過剰に含有すると、粗大な酸化物を形成して延性を極端に劣化させるため、その量を０．００１〜０．１％とした。 Al is an element effective for deoxidation, austenite grain size reduction, etc., and in order to exhibit 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, so the amount was made 0.001 to 0.1%.
Ｎは、ＡｌやＴｉと化合してオーステナイト粒微細化に有効に働くため、微量であれば機械的性質の向上に寄与する。また、工業的に鋼中のＮを完全に除去することは不可能であり、必要以上に低減することは製造工程に過大な負担をかけるため好ましくない。そのため工業的に制御が可能で、製造工程への負荷が許容できる範囲として下限を０．００１％とする。過剰に含有すると固溶Ｎが増加し、歪時効特性が劣化するために、上限を０．００８％とした。 Since N combines with Al and Ti and effectively works to refine the austenite grains, it contributes to the improvement of the mechanical properties if the amount is small. Moreover, it is impossible to remove N in steel completely industrially, and reducing it more than necessary is not preferable because it places an excessive burden on the manufacturing process. Therefore, it is industrially controllable and the lower limit is set to 0.001% as a range in which the load on the manufacturing process can be tolerated. If excessively contained, the solid solution N increases and the strain aging characteristics deteriorate, so the upper limit was made 0.008%.
以上が本発明の基本成分の限定理由であるが、本発明においては、強度・靭性の調整のために、必要に応じてＣｕ、Ｎｉ、Ｃｒ、Ｍｏ、Ｗ、Ｔｉ、Ｎｂ、Ｚｒ、Ｖ、Ｂの１種あるいは２種以上を含有することができる。以下に各元素の成分限定理由を述べる。 The above is the reason for limiting the basic components of the present invention. In the present invention, Cu, Ni, Cr, Mo, W, Ti, Nb, Zr, V, One or more of B can be contained. The reasons for limiting the components of each element will be described below.
Ｃｕは、靭性を低下させずに強度の上昇に有効な元素であるが、０．１％未満では効果がなく、２．５％を超えると鋼片加熱時や溶接時に割れを生じやすくする。従ってその量を０．１〜２．５％とした。 Cu is an element effective for increasing the strength without reducing toughness, but if it is less than 0.1%, there is no effect, and if it exceeds 2.5%, it tends to cause cracks during heating of the steel slab or during welding. Therefore, the amount is set to 0.1 to 2.5%.
Ｎｉは、靭性および強度の改善に有効な元素であり、その効果を得るためには０．１％以上の添加が必要であるが、５％を超える過剰な添加では、効果が飽和する一方で、ＨＡＺ靭性や溶接性の劣化を生じる懸念があり、また高価な元素であるため経済性も考慮して、その量を０．１〜５％とした。 Ni is an element effective for improving toughness and strength. To obtain the effect, Ni is required to be added in an amount of 0.1% or more. However, excessive addition exceeding 5% saturates the effect. There is a concern that the HAZ toughness and weldability may be deteriorated, and since it is an expensive element, its amount is set to 0.1 to 5% in consideration of economy.
Ｃｒは、焼入れ性を高めて強度を確保する上で０．０１％以上必要である。一方、１．５％を超えるとＮｉと同様の理由で好ましくない。従ってその量を０．０１〜１．５％とした。 Cr is required to be 0.01% or more for enhancing the hardenability and ensuring the strength. On the other hand, if it exceeds 1.5%, it is not preferable for the same reason as Ni. Therefore, the amount is set to 0.01 to 1.5%.
Ｍｏは、焼入れ性向上、強度向上、耐焼戻し脆化、および再結晶抑制に有効な元素で、その効果を得るためには０．０１％以上の添加が必要であるが、１．５％を超えると靭性および溶接性が劣化する。従ってその量を０．０１〜１．５％とした。 Mo is an element effective for improving hardenability, improving strength, resistance to temper embrittlement, and suppressing recrystallization. To obtain the effect, addition of 0.01% or more is necessary. If exceeded, toughness and weldability deteriorate. Therefore, the amount is set to 0.01 to 1.5%.
Ｗは、焼入れ性を高めて強度を確保するのに必要な元素であるが、効果を発揮でき、他特性に悪影響を及ぼさない範囲として、その量を０．０１〜１．５％とした。 W is an element necessary for enhancing the hardenability and securing the strength, but the amount is set to 0.01 to 1.5% as a range that can exert the effect and does not adversely affect other characteristics.
Ｔｉは、析出強化により母材強度向上に寄与するとともに、高温でも安定なＴｉＮの形成により加熱オーステナイト粒径微細化にも有効な元素であり、効果を発揮するためには０．００１％以上含有する必要がある。一方、０．０５％を超えると粗大な酸化物を形成して延性を極端に劣化させるため、その量を０．００１〜０．０５％とした。 Ti is an element that contributes to improving the strength of the base metal by precipitation strengthening, and is also an effective element for refining the heated austenite grain size by forming TiN that is stable even at high temperatures. There is a need to. On the other hand, if it exceeds 0.05%, a coarse oxide is formed and the ductility is extremely deteriorated. Therefore, the amount is set to 0.001 to 0.05%.
Ｎｂ、Ｚｒ、Ｖは、析出強化により母材の強度向上に寄与するが、０．００５％未満では効果がなく、０．２％を超える過剰の添加では延性や靭性が劣化する。従って、Ｎｂ、Ｚｒ、Ｖともにその量を０．００５〜０．２％とした。 Nb, Zr, and V contribute to improving the strength of the base metal by precipitation strengthening, but if less than 0.005%, there is no effect, and if it exceeds 0.2%, ductility and toughness deteriorate. Therefore, the amount of Nb, Zr, and V is set to 0.005 to 0.2%.
Ｂは、固溶状態でオーステナイト粒界に偏析することで、微量で焼入れ性を高めることが可能な元素であるが、粒界に偏析した状態ではオーステナイトの再結晶抑制にも有効である。焼入れ性、再結晶抑制に効果を発揮するためには０．０００２％以上の添加が必要であるが、一方、０．００５％を超える過剰の添加は、粗大な析出物を生じて靭性が劣化するため、その量を０．０００２〜０．００５％とした。 B is an element capable of improving the hardenability in a small amount by segregating at the austenite grain boundary in a solid solution state. However, B is also effective in suppressing recrystallization of austenite in the state segregated at the grain boundary. Addition of 0.0002% or more is necessary to exert effects on hardenability and recrystallization suppression. On the other hand, excessive addition exceeding 0.005% produces coarse precipitates and deteriorates toughness. Therefore, the amount was made 0.0002 to 0.005%.
さらに、本発明においては、延性の向上、継手靭性の向上のために、必要に応じてＭｇ、Ｃａ、ＲＥＭの１種または２種以上を添加することができる。
Ｍｇ、Ｃａ、ＲＥＭはいずれも硫化物の熱間圧延中の展伸を抑制して延性向上に有効である。酸化物を微細化させて継手靭性の向上にも有効に働く。その効果を発揮するための下限の含有量は、Ｍｇは０．０００５％、Ｃａは０．０００５％、ＲＥＭは０．００５％である。一方、過剰に含有すると、硫化物や酸化物の粗大化を生じ、延性、靭性の劣化を招くため、上限の含有量を、Ｍｇは０．０１％、Ｃａは０．０１％、ＲＥＭは０．０５％とした。
Furthermore, in this invention, 1 type, or 2 or more types of Mg, Ca, and REM can be added as needed for the improvement of ductility and the improvement of joint toughness.
Mg, Ca, and REM are all effective in improving ductility by suppressing extension during hot rolling of sulfides. 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 and toughness. Therefore, the upper content is 0.01% for Mg, 0.01% for Ca, and 0 for REM. .05%.
以上が、本発明の基本要件であるミクロ組織と化学成分の限定理由である。加えて、本発明の組織要件を満足させるための適切な製造方法についても提示する。この製造方法は、適正な化学成分を含有する鋼片を、Ａｃ3変態点以上１３５０℃以下の温度に加熱後、Ａｒ3変態点〜１２５０℃のオーステナイト単相域で累積圧下率が５３〜８０％で圧延した後、圧延開始温度がＡｒ3変態点以下、圧延終了温度が６００℃以上のオーステナイト−フェライトの二相域で、累積圧下率が４０〜９０％の仕上圧延を行う。仕上圧延後、５〜８０℃／ｓの冷却速度で２０〜４００℃まで加速冷却してもよい。さらに、３００〜５００℃の温度範囲で焼戻ししてもよい。ただし、本発明のミクロ組織については、その達成手段を問わず効果を発揮するものである。 The above is the reason for limiting the microstructure and chemical components, which are the basic requirements of the present invention. In addition, a suitable manufacturing method for satisfying the organizational requirements of the present invention is also presented. In this production method, a steel slab containing an appropriate chemical component is heated to a temperature not lower than the Ac3 transformation point and not higher than 1350 ° C, and then the cumulative reduction ratio is 53 to 80% in the austenite single phase region of Ar3 transformation point to 1250 ° C. After rolling, finish rolling is performed in an austenite-ferrite two-phase region in which the rolling start temperature is equal to or lower than the Ar3 transformation point and the rolling end temperature is 600 ° C. or higher, and the cumulative rolling reduction is 40 to 90%. After finish rolling, accelerated cooling to 20 to 400 ° C. may be performed at a cooling rate of 5 to 80 ° C./s. Furthermore, you may temper in the temperature range of 300-500 degreeC. However, the microstructure of the present invention, Ru der which effective regardless of its achievement means.
熱間圧延に先立ち、鋼塊を１００％オーステナイト化する必要があり、このためには鋼塊の温度をＡｃ3変態点以上に加熱する必要がある。しかし、１３５０℃を超えて加熱すると、オーステナイト粒が著しく粗大化し、圧延後に細粒フェライトが得られなくなるので、加熱温度の上限は１３５０℃とする。
Prior to hot rolling, it is necessary to make the
引き続く熱間圧延をＡｒ3変態点〜１２５０℃の温度域に限定したのは、オーステナイト単相域での圧延を施すことによって、変態温度の高温化と変態組織の微細化が図られ、二相域圧延において細粒フェライトが得られるからである。累積圧下率１０％未満ではこの効果は少ないが、下限は、実施例の表２に示す本発明例Ａ２の累積圧下率に基づき５３％とするのが適切である。また８０％を超えると、引き続く二相域圧延での圧下を確保できなくなる。この場合、オーステナイト域で制御圧延を施し、二相域圧延の前にオーステナイト粒をさらに微細化しておく方が好ましい。 The subsequent hot rolling was limited to the temperature range of Ar3 transformation point to 1250 ° C. By rolling in the austenite single phase region, the transformation temperature was increased and the transformation structure was refined. This is because fine-grained ferrite can be obtained by rolling. This effect is cumulative rolling reduction below 10 percent less Iga, lower limit, it is appropriate to 53% based on the cumulative rolling reduction of the Invention Example A2 shown in Table 2 of Example. On the other hand, if it exceeds 80%, it becomes impossible to secure the reduction in the subsequent two-phase rolling . In this case, it is preferable to perform controlled rolling in the austenite region and further refine the austenite grains before the two-phase region rolling.
本発明では、硬いフェライト中に扁平で硬いマルテンサイトを層状に分散させることが必要であり、このためにＡｒ3変態点以下における仕上圧延が極めて重要な役割を果たし、本発明で必須の工程である。フェライトの硬さ向上、マルテンサイトの硬さ向上、扁平化、変態開始温度の低温化のためには、Ａｒ3点以下の仕上圧延が必要であり、圧延温度は低いほうが望ましいが、低温ほど変形抵抗が上昇するので圧延荷重が上昇し、圧延が困難である。また６００℃未満になると、マルテンサイトの分率を５％以上確保することができなくなる。従って圧延終了温度を６００℃以上とした。 In the present invention, it is necessary to disperse flat and hard martensite in a layer form in hard ferrite, and for this reason, finish rolling below the Ar3 transformation point plays an extremely important role and is an essential step in the present invention. . In order to improve the hardness of ferrite, the hardness of martensite, flattening, and lowering the transformation start temperature, finishing rolling at Ar3 point or lower is necessary, and the lower the rolling temperature is desirable, the lower the deformation resistance. As a result, the rolling load increases and rolling is difficult. On the other hand, when the temperature is lower than 600 ° C., it is impossible to secure a martensite fraction of 5% or more. Therefore, the rolling end temperature is set to 600 ° C. or higher.
仕上圧延の累積圧下率が４０％未満では、フェライトの硬さ向上、マルテンサイトの硬さ向上、扁平化の効果が少なく、さらにフェライトとマルテンサイトの板厚方向の層間隔が増大してしまうため、累積圧下率は大きいほど好ましい。従って仕上圧延累積圧下率を４０〜９０％とした。 If the cumulative rolling reduction of finish rolling is less than 40%, the effect of improving the hardness of the ferrite, improving the hardness of the martensite, and flattening is small, and further, the layer spacing in the thickness direction of the ferrite and martensite increases. The larger the cumulative rolling reduction, the better. Therefore, the finish rolling cumulative reduction ratio is set to 40 to 90%.
二相域圧延後の冷却方法としては、マルテンサイト変態させるために、マルテンサイト変態開始温度以下まで、５〜８０℃／ｓの冷却速度で２０〜４００℃まで加速冷却する必要がある。
加速冷却する場合の冷却速度を５〜８０℃／ｓに限定したのは、５℃／ｓ未満では加速冷却にマルテンサイト変態が困難である上に、母材の強度、靭性の向上が期待できないためであり、８０℃／ｓ超では表層と内部との組織あるいは特性の差が大きく生じて好ましくないためである。
As a cooling method after the two-phase region rolling, in order to perform martensite transformation, it is necessary to accelerate cooling to 20 to 400 ° C. at a cooling rate of 5 to 80 ° C./s up to the martensite transformation start temperature or lower.
The reason why the cooling rate in the case of accelerated cooling is limited to 5 to 80 ° C./s is that if it is less than 5 ° C./s, martensitic transformation is difficult for accelerated cooling, and improvement in strength and toughness of the base material cannot be expected. This is because if the temperature exceeds 80 ° C./s, a difference in structure or characteristics between the surface layer and the inside greatly occurs, which is not preferable.
また、加速冷却は鋼板の所望の強度、靭性レベルに応じて２０〜４００℃で停止する。加速冷却の停止温度を２０℃未満とすることは材質を制御する上でなんら効果がなく、単に製造コストの上昇を招くだけで意味がない。逆に加速冷却を４００℃超で停止すると、マルテンサイト変態が困難である上に、内部応力が緩和され、疲労き裂伝播特性の向上が期待できない。 The accelerated cooling stops at 20 to 400 ° C. depending on the desired strength and toughness level of the steel sheet. Setting the stop temperature of accelerated cooling to less than 20 ° C. has no effect in controlling the material, and merely causes an increase in manufacturing cost and is meaningless. Conversely, if accelerated cooling is stopped at over 400 ° C., martensitic transformation is difficult, internal stress is relaxed, and fatigue crack propagation characteristics cannot be expected.
圧延・冷却後に、必要に応じて引き続き実施する焼戻し処理は、回復による母材組織の靭性向上を目的としたものであるから、加熱温度は逆変態が生じない温度域であるＡｃ1以下でなければならない。さらに５００℃超では、内部応力が緩和されることにより疲労き裂伝播特性が劣化することから上限を５００℃とした。また、回復は転位の消滅・合体により格子欠陥密度を減少させるものであり、これを実現させるためには３００℃以上に加熱することが必要であるため、下限を３００℃とした。なお、この焼戻し熱処理によって、生成する焼戻しマルテンサイトも本発明の組織要件であるマルテンサイトとして定義する。 The tempering process that is subsequently carried out as necessary after rolling and cooling is intended to improve the toughness of the base metal structure by recovery. Therefore, the heating temperature must be less than Ac1, which is a temperature range in which reverse transformation does not occur. Don't be. Further, if the temperature exceeds 500 ° C., the fatigue crack propagation characteristics deteriorate due to relaxation of internal stress, so the upper limit was set to 500 ° C. Further, the recovery is to reduce the lattice defect density by the disappearance and coalescence of dislocations, and in order to realize this, it is necessary to heat to 300 ° C. or higher, so the lower limit was set to 300 ° C. In addition, the tempered martensite produced | generated by this tempering heat processing is also defined as a martensite which is the structure | tissue requirements of this invention.
以下に、本発明の効果を実施例によってさらに具体的に述べる。
実施例に用いた供試鋼の化学成分を表１に示す。各供試鋼は造塊後、分塊圧延により、あるいは連続鋳造により鋼片としたものである。表１の鋼番１〜２０は本発明の化学組成範囲を満足しており、鋼番２１〜２５は本発明の化学組成範囲を満足していない。
Hereinafter, the effects of the present invention will be described more specifically with reference to examples.
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. Steel numbers 1 to 20 in Table 1 satisfy the chemical composition range of the present invention, and steel numbers 21 to 25 do not satisfy the chemical composition range of the present invention.
表１の化学成分の鋼片を表２に示す条件により鋼板に製造した。試験Ｎｏ．Ａ１〜Ａ２３は請求項４〜６に関連した方法により製造した。また、試験Ｎｏ．Ｂ１〜Ｂ１２は本発明の製造条件を満足していない。それぞれの室温での機械的性質を表２に合わせて示す。
表３は前述の鋼番１〜２５、試験Ｎｏ．Ａ１〜Ａ２３、Ｂ１〜Ｂ１２からなる鋼板のミクロ組織調査結果、疲労試験結果を示す。
Steel strips having the chemical components shown in Table 1 were produced into steel plates under the conditions shown in Table 2. Test No. A1 to A23 were produced by the method according to claims 4-6. In addition, Test No. B1 to B12 do not satisfy the production conditions of the present invention. Table 2 shows the mechanical properties of each at room temperature.
Table 3 shows steel Nos. 1 to 25, test Nos. The microstructure check result of the steel plate which consists of A1-A23 and B1-B12 and the fatigue test result are shown.
ミクロ組織は、鋼板の圧延方向の板厚断面を鏡面研磨後、ナイタール腐食、レペラ腐食によって現出させ、光学顕微鏡を用いて観察し、生成した相を後述の硬さ試験結果と併用して同定した。そして、硬さ測定はマイクロビッカース硬さ試験機を用いて荷重１０ｇにて実施した。各相の分率、アスペクト比、層間隔は、光学顕微鏡写真を画像解析することによって求めた。 The microstructure is obtained by mirror-polishing the thickness cross section in the rolling direction of the steel sheet, revealing it by nital corrosion and repeller corrosion, observing it with an optical microscope, and identifying the generated phase in combination with the hardness test results described below. did. The hardness was measured using a micro Vickers hardness tester with a load of 10 g. The fraction, aspect ratio, and layer spacing of each phase were determined by image analysis of optical micrographs.
図１は、疲労き裂伝播試験に用いた試験片を示す図である。疲労き裂伝播試験条件は以下の通りとした。
・荷重負荷方式：３点曲げ、
・応力比 ：０．１、
・環境 ：室温大気中、
・き裂長さ測定：直流電位差法
FIG. 1 is a view showing a test piece used in a fatigue crack propagation test. The fatigue crack propagation test conditions were as follows.
・ Loading method: 3-point bending,
-Stress ratio: 0.1,
・ Environment: In room temperature atmosphere,
・ Crack length measurement: DC potential difference method
図２は、溶接継手疲労試験に用いた試験片を示す図である。溶接は、入熱１８ｋＪ／ｍｉｎで炭酸ガスアーク溶接を行った。疲労試験条件は以下の通りとした。
・荷重負荷方式：軸力、
・応力比 ：０．１、
・環境 ：室温大気中、
・試験応力範囲：１５０ＭＰａ
FIG. 2 is a view showing a test piece used in a welded joint fatigue test. For welding, carbon dioxide arc welding was performed at a heat input of 18 kJ / min. The fatigue test conditions were as follows.
・ Loading method: axial force,
-Stress ratio: 0.1,
・ Environment: In room temperature atmosphere,
Test stress range: 150 MPa
試験Ｎｏ．Ａ１〜Ａ２０は、いずれも本発明の化学組成の鋼片を本発明の要件に従って製造した鋼材であり、組織要件も満足しており、応力拡大係数範囲ΔＫが２０ＭＰａ√ｍのときの疲労き裂伝播速度ｄａ／ｄＮが１０−８ｍ／ｃｙｃｌｅ以下、かつ溶接継手疲労寿命が試験Ｎｏ．Ｂ１の比較例に対して２倍以上と、優れた疲労特性を有していた。 Test No. Each of A1 to A20 is a steel material produced by producing a steel slab having the chemical composition of the present invention in accordance with the requirements of the present invention, which also satisfies the structural requirements, and a fatigue crack when the stress intensity factor range ΔK is 20 MPa√m. Propagation velocity da / dN is 10 −8 m / cycle or less, and the weld joint fatigue life is Test No. As compared with the comparative example of B1, the fatigue characteristics were excellent, twice or more.
一方、試験Ｎｏ．Ａ２１〜２３は、本発明の製造要件は満足しているが、化学組成の限定範囲が外れている。試験Ｎｏ．Ａ２１、２３は、フェライト−マルテンサイト組織となっているが、マルテンサイト分率が小さいか、または層間隔が大きいため、ΔＫ＝２０ＭＰａ√ｍのときの伝播速度が１０−８ｍ／ｃｙｃｌｅ以上であり、そのため溶接継手疲労寿命が試験Ｎｏ．Ｂ１の比較例に対して２倍以下であり、本発明鋼に比べて疲労特性は劣っていた。また試験Ｎｏ．Ａ２２は、マルテンサイト分率が過剰であるため、靭性が大幅に劣化し、溶接継手疲労寿命は疲労試験途中で脆性破壊を起こしたため、本発明鋼に比べて著しく劣っていた。また、層間隔が小さすぎたことにより伝播特性も本開発鋼に比べ劣っていた。 On the other hand, test no. In A21 to 23, the production requirements of the present invention are satisfied, but the chemical composition is not within the limited range. Test No. A21 and 23 have a ferrite-martensite structure, but since the martensite fraction is small or the layer spacing is large, the propagation speed when ΔK = 20 MPa√m is 10 −8 m / cycle or more. Therefore, the weld joint fatigue life is Test No. The fatigue property was inferior to that of the comparative steel of B1, and the fatigue properties were inferior to the steel of the present invention. In addition, Test No. Since A22 has an excessive martensite fraction, the toughness was greatly deteriorated, and the fatigue life of the welded joint was significantly inferior to that of the steel of the present invention because brittle fracture occurred during the fatigue test. In addition, the propagation characteristics were inferior to the newly developed steel due to the too small layer spacing.
また試験Ｎｏ．Ｂ１〜Ｂ１０は、本発明の化学組成の限定範囲は満足しているが、製造要件が外れている。試験Ｎｏ．Ｂ１、Ｂ６、Ｂ７、Ｂ８、Ｂ１０は、第二相がマルテンサイトではなく、マルテンサイト以外ではき裂進展の有効な障害とはなり難いため、疲労き裂伝播特性は本発明鋼に比べ劣化し、溶接継手疲労寿命も向上しなかった。
試験Ｎｏ．Ｂ２、Ｂ３は、第二相がマルテンサイトであるが、フェライト硬さが小さく、内部応力が高められなかったことに加え、アスペクト比が小さいか、または層間隔が大きかったので、き裂進展時にマルテンサイトに当たる頻度が少なく、有効な障害となり得なかったため、疲労き裂伝播特性は本発明鋼に比べ劣化し、溶接継手疲労寿命も向上しなかった。
In addition, Test No. Although B1-B10 is satisfying the limited range of the chemical composition of the present invention, the production requirements are deviated. Test No. In B1, B6, B7, B8, and B10, the second phase is not martensite, and other than martensite is less likely to be an effective obstacle to crack growth. Therefore, fatigue crack propagation characteristics deteriorate compared to the steel of the present invention. The fatigue life of welded joints was not improved.
Test No. In B2 and B3, the second phase is martensite, but the ferrite hardness is small and the internal stress cannot be increased. In addition, the aspect ratio is small or the layer spacing is large. Since the frequency of hitting martensite was low and could not be an effective obstacle, the fatigue crack propagation characteristics were deteriorated compared to the steel of the present invention, and the fatigue life of the welded joint was not improved.
試験Ｎｏ．Ｂ４は、焼戻し温度が高く、内部応力が緩和されたため、き裂進展の障害とはならず、疲労特性は本発明鋼に比べ劣っていた。試験Ｎｏ．Ｂ５は仕上げ圧延開始温度、加速冷却開始温度が高く、第二相の大部分がベイナイトとなったため、き裂進展の障害とはならず、疲労特性は本発明鋼に比べ劣っていた。試験Ｎｏ．Ｂ９は、第二相がマルテンサイトであるが、仕上げ累積圧下率が小さく、アスペクト比が極めて小さかったため、き裂進展時にマルテンサイトに当たる頻度が少なく、有効な障害となり得なかったため、疲労特性は本発明鋼に比べ劣っていた。
さらに、試験Ｎｏ．Ｂ１１〜Ｂ１２については、化学組成、製造法ともに本発明の限定範囲を満たしてないため、本発明鋼に比べて疲労特性は顕著に劣化していた。
Test No. Since B4 had a high tempering temperature and the internal stress was relaxed, B4 did not become an obstacle to crack growth, and the fatigue properties were inferior to the steel of the present invention. Test No. B5 had a high finish rolling start temperature and an accelerated cooling start temperature, and most of the second phase became bainite, so it did not become an obstacle to crack growth and the fatigue properties were inferior to the steel of the present invention. Test No. In B9, the second phase is martensite, but since the finish cumulative reduction ratio was small and the aspect ratio was very small, the frequency of hitting martensite during crack growth was low and could not be an effective failure. It was inferior to the inventive steel.
Furthermore, test no. Regarding B11 to B12, both the chemical composition and the production method did not satisfy the limited range of the present invention, so the fatigue characteristics were significantly deteriorated compared to the steel of the present invention.
Claims (3)
Ｃ ：０．０５〜０．２％、
Ｓｉ：０．０１〜０．８％、
Ｍｎ：０．５〜２％、
Ｐ ：０．０２％以下、
Ｓ ：０．０２％以下、
Ａｌ：０．００１〜０．１％、
Ｎ ：０．００１〜０．００８％
を含有し、残部Ｆｅおよび不可避不純物からなり、母材のミクロ組織が、ビッカース硬さが１５０以上のフェライトを母相とし、ビッカース硬さが４００〜９００、面積率が５〜３０％、アスペクト比（長軸／短軸）が３以上の扁平なマルテンサイトを第二相とした層状組織であり、さらにフェライトとマルテンサイトの板厚方向の平均層間隔が３〜５０μｍであり、応力比０．１の応力拡大係数範囲ΔＫが２０ＭＰａ√ｍのときの疲労き裂伝播速度ｄａ／ｄＮが１０−８ｍ／ｃｙｃｌｅ以下であることを特徴とする、疲労き裂伝播特性に優れた鋼板。 % By mass
C: 0.0 5 ~0.2%,
Si: 0.01~ 0.8%,
Mn: 0.5-2%
P: 0.02% or less,
S: 0.02% or less,
Al: 0.001 to 0.1%,
N: 0.001 to 0.008%
Containing the balance Fe and unavoidable impurities, and the microstructure of the base material is a ferrite having a Vickers hardness of 150 or more, the Vickers hardness is 400 to 900, the area ratio is 5 to 30%, and the aspect ratio (Major axis / minor axis) is a layered structure having flat martensite of 3 or more as a second phase, and the average layer spacing in the plate thickness direction of ferrite and martensite is 3 to 50 μm, and the stress ratio is 0. A steel plate having excellent fatigue crack propagation characteristics, wherein a fatigue crack propagation rate da / dN when a stress intensity factor range ΔK of 1 is 20 MPa√m is 10 −8 m / cycle or less.
Ｃｕ：０．１〜２．５％、
Ｎｉ：０．１〜５％、
Ｃｒ：０．０１〜１．５％、
Ｍｏ：０．０１〜１．５％、
Ｗ ：０．０１〜１．５％、
Ｔｉ：０．００１〜０．０５％、
Ｎｂ：０．００５〜０．２％、
Ｚｒ：０．００５〜０．２％
Ｖ ：０．００５〜０．２％、
Ｂ ：０．０００２〜０．００５％
の１種または２種以上を含有することを特徴とする、請求項１に記載の疲労き裂伝播特性に優れた鋼板。 In addition,
Cu: 0.1 to 2.5%,
Ni: 0.1 to 5%,
Cr: 0.01 to 1.5%
Mo: 0.01 to 1.5%,
W: 0.01 to 1.5%,
Ti: 0.001 to 0.05%,
Nb: 0.005 to 0.2%,
Zr: 0.005 to 0.2%
V: 0.005 to 0.2%,
B: 0.0002 to 0.005%
The steel plate excellent in fatigue crack propagation characteristics according to claim 1, comprising one or more of the following.
Ｍｇ：０．０００５〜０．０１％、
Ｃａ：０．０００５〜０．０１％、
ＲＥＭ：０．００５〜０．０５％
の１種または２種以上を含有することを特徴とする、請求項１または２に記載の疲労き裂伝播特性に優れた鋼板。 In addition,
Mg: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
REM: 0.005 to 0.05%
The steel plate excellent in fatigue crack propagation characteristics according to claim 1 or 2, characterized by containing one or more of the following.
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| JP2005016036A JP4926406B2 (en) | 2004-04-08 | 2005-01-24 | Steel sheet with excellent fatigue crack propagation characteristics |
| PCT/JP2005/007189 WO2005098069A1 (en) | 2004-04-08 | 2005-04-07 | Steel plate excellent in fatigue crack propagation characteristics and method for production thereof |
| KR1020067023289A KR20060130779A (en) | 2004-04-08 | 2005-04-07 | Steel plate excellent in fatigue crack propagation characteristics and method for production thereof |
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| JP4998708B2 (en) * | 2007-02-26 | 2012-08-15 | Ｊｆｅスチール株式会社 | Steel material with small material anisotropy and excellent fatigue crack propagation characteristics and method for producing the same |
| JP5407143B2 (en) * | 2007-03-09 | 2014-02-05 | Ｊｆｅスチール株式会社 | Fatigue crack propagation retarding steel and its manufacturing method |
| JP5407144B2 (en) * | 2007-03-09 | 2014-02-05 | Ｊｆｅスチール株式会社 | Steel material with excellent fatigue crack growth control |
| JP5037203B2 (en) * | 2007-04-12 | 2012-09-26 | 新日本製鐵株式会社 | Method for producing high-strength steel material having yield stress of 470 MPa or more and tensile strength of 570 MPa or more excellent in toughness of weld heat-affected zone |
| JP5037204B2 (en) * | 2007-04-12 | 2012-09-26 | 新日本製鐵株式会社 | Method for producing high-strength steel material having yield stress of 500 MPa or more and tensile strength of 570 MPa or more which is excellent in toughness of weld heat affected zone |
| KR100957963B1 (en) * | 2007-12-26 | 2010-05-17 | 주식회사 포스코 | Steel for a structure having excellent low temperature toughnetss, tensile strength and low yield ratio, of heat affected zone and manufacturing method for the same |
| TWI341332B (en) * | 2008-01-07 | 2011-05-01 | Nippon Steel Corp | Wear-resistant steel sheet having excellent wear resistnace at high temperatures and excellent bending workability and method for manufacturing the same |
| JP5284015B2 (en) * | 2008-09-04 | 2013-09-11 | 株式会社神戸製鋼所 | Thick steel plate with excellent brittle crack propagation stop properties |
| KR101091306B1 (en) * | 2008-12-26 | 2011-12-07 | 주식회사 포스코 | High Strength Steel Plate for Containment Vessel of Atomic Plant and Manufacturing Method Thereof |
| KR101149258B1 (en) * | 2009-04-27 | 2012-05-25 | 현대제철 주식회사 | Steel with superior anisotropy for impact toughness and the method of producing the same |
| KR101271990B1 (en) * | 2009-12-01 | 2013-06-05 | 주식회사 포스코 | High strength steel sheet and method for manufacturing the same |
| KR100981856B1 (en) * | 2010-02-26 | 2010-09-13 | 현대하이스코 주식회사 | Method of manufacturing high strength steel sheet with excellent coating characteristics |
| CN102091893A (en) * | 2010-12-30 | 2011-06-15 | 哈尔滨工业大学 | Design method capable of ensuring welding joint to be born according to bearing capability of parent metal |
| JP6228491B2 (en) * | 2013-08-26 | 2017-11-08 | 株式会社神戸製鋼所 | Thick steel plate with excellent fatigue characteristics and method for producing the same |
| JP6064897B2 (en) * | 2013-12-27 | 2017-01-25 | Ｊｆｅスチール株式会社 | High-strength steel material with excellent fatigue crack propagation resistance and its determination method |
| JP6303782B2 (en) * | 2014-05-08 | 2018-04-04 | 新日鐵住金株式会社 | Hot-rolled steel sheet and manufacturing method thereof |
| PL3162908T3 (en) | 2014-07-14 | 2019-10-31 | Nippon Steel & Sumitomo Metal Corp | Hot-rolled steel sheet |
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