JP2013104070A - High-strength steel excellent in delayed breakage resistance, and high-strength bolt - Google Patents
High-strength steel excellent in delayed breakage resistance, and high-strength bolt Download PDFInfo
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本発明は、耐遅れ破壊特性に優れた鋼に関し、例えばボルト用鋼を代表とする、引張強さ1200MPa超の耐遅れ破壊特性に優れた高強度鋼および、その高強度鋼からなる高強度ボルトに関するものである。 The present invention relates to a steel excellent in delayed fracture resistance, for example, a high strength steel excellent in delayed fracture resistance with a tensile strength exceeding 1200 MPa, such as a steel for bolts, and a high strength bolt made of the high strength steel. It is about.
自動車、産業機械、橋梁、土木建築等、各種産業分野で使用されている高強度鋼は、例えばJIS G4053に規定されるSCr(クロム鋼)、SCM(クロムモリブデン鋼)等の機械構造用合金鋼であり、その鋼に焼入れ、焼戻し処理を施すことによって製造されている。しかし、上記の鋼は、引張強さ1200MPaを超えると耐遅れ破壊特性が著しく低下し、使用中に環境から侵入する水素に起因する遅れ破壊を生じる危険性が増大することはよく知られている。そのため、例えば、土木建築向けの鋼では、規格は引張強さ1200MPa以下であり、実用上、引張強さが1150MPa級の鋼に制限されている。 High-strength steel used in various industrial fields such as automobiles, industrial machinery, bridges, civil engineering and construction is an alloy steel for machine structures such as SCr (chromium steel) and SCM (chromium-molybdenum steel) defined in JIS G4053. It is manufactured by quenching and tempering the steel. However, it is well known that when the above steel strength exceeds 1200 MPa, the delayed fracture resistance significantly decreases and the risk of delayed fracture due to hydrogen entering from the environment during use increases. . Therefore, for example, in steel for civil engineering and construction, the standard is a tensile strength of 1200 MPa or less, and the tensile strength is practically limited to steel of 1150 MPa class.
鋼材の高強度化にともない、耐遅れ破壊特性を改善した高強度鋼及びその製造方法が種々提案されている。例えば、特許文献1には、旧オーステナイト粒を微細化させることを図ることによって、耐遅れ破壊特性を改善した発明が記載されている。旧オーステナイト粒を微細化する方法は、耐遅れ破壊特性の改善効果は認められるものの、大幅な改善には至っていない。 With increasing strength of steel materials, various high strength steels with improved delayed fracture resistance and methods for producing the same have been proposed. For example, Patent Document 1 describes an invention in which delayed fracture resistance is improved by refining prior austenite grains. Although the method of refining prior austenite grains has been confirmed to have an effect of improving delayed fracture resistance, it has not yet improved significantly.
また、例えば、特許文献2、特許文献3、および特許文献4には、鋼中に水素をトラップさせる酸化物、炭化物、窒化物の単独あるいは複合析出物を分散分布させることにより、遅れ破壊が発現する臨界の水素量を増加させることにより耐遅れ破壊特性を改善した発明が記載されている。これら発明において、耐遅れ破壊特性を改善する機構の一つに焼入れ、焼戻し処理で生成する析出物を活用する技術思想が提案されている。この耐遅れ破壊特性を改善するためには、水素トラップさせる析出物を最適に分散分布させる鋼の化学成分や熱処理条件の限定が必須である。 Further, for example, in Patent Document 2, Patent Document 3, and Patent Document 4, delayed fracture occurs by dispersing or distributing oxides, carbides, nitrides alone or composite precipitates that trap hydrogen in steel. The invention describes improved delayed fracture resistance by increasing the critical hydrogen content. In these inventions, a technical idea of utilizing precipitates generated by quenching and tempering has been proposed as one of the mechanisms for improving delayed fracture resistance. In order to improve the delayed fracture resistance, it is essential to limit the chemical composition of the steel and the heat treatment conditions that optimally distribute and distribute the precipitate to be trapped with hydrogen.
本発明は、鋼材の高強度化にともない現出する耐遅れ破壊現象に代表される水素脆化をより抑制することのできる、耐遅れ破壊特性に優れた高強度鋼、および、高強度ボルトを提供することを目的とする。 The present invention provides a high-strength steel excellent in delayed fracture resistance and a high-strength bolt capable of further suppressing hydrogen embrittlement represented by the delayed fracture resistance phenomenon that appears as the strength of steel materials increases. The purpose is to provide.
本発明者らは、遅れ破壊特性に及ぼす各種因子について鋭意検討し、以下の知見を見出した。
(a)鋼中に多量のFe系炭化物εを析出することによって、遅れ破壊が発現する臨界の水素量(以下、「限界拡散性水素量」と記載する)は増加し、耐遅れ破壊特性が向上する。これは、Fe系炭化物εが外部環境から侵入した鋼中の水素をトラップし、限界拡散性水素量を増加させるためである。
(b)Siを多量に添加することによって、Fe系炭化物εが安定化し、より高温で焼戻し処理が可能となるため、母相の転位等の欠陥密度が減少し水素に対する脆化感受性が低下する。
The present inventors diligently studied various factors affecting delayed fracture characteristics and found the following findings.
(A) By depositing a large amount of Fe-based carbide ε in the steel, the critical hydrogen amount at which delayed fracture occurs (hereinafter referred to as “critical diffusible hydrogen amount”) is increased, and delayed fracture resistance is improved. improves. This is because the Fe-based carbide ε traps hydrogen in the steel that has entered from the outside environment and increases the amount of critical diffusible hydrogen.
(B) By adding a large amount of Si, Fe-based carbide ε is stabilized and tempering treatment is possible at a higher temperature, so that the defect density such as dislocations in the parent phase is reduced and the embrittlement sensitivity to hydrogen is reduced. .
本発明は上記知見に基づいて完成したもので、その発明の要旨とするところは、次の通りである。 The present invention has been completed based on the above findings, and the gist of the present invention is as follows.
(1)質量%で、
C:0.20〜1.50%、
Si:0.20〜5.00%、
Mn:0.10〜3.00%、
P:0.0005〜0.1000%、
S:0.0005〜0.2000%、
N:0.0020〜0.0200%
を含有し、残部がFe及び不可避的不純物からなり、鋼組織はFe系炭化物εが微細分散した焼戻しマルテンサイト組織であることを特徴とする耐遅れ破壊特性に優れた高強度鋼。
(1) In mass%,
C: 0.20 to 1.50%,
Si: 0.20 to 5.00%,
Mn: 0.10 to 3.00%,
P: 0.0005 to 0.1000%,
S: 0.0005 to 0.2000%,
N: 0.0020 to 0.0200%
A high-strength steel excellent in delayed fracture resistance, characterized in that the balance is made of Fe and inevitable impurities, and the steel structure is a tempered martensite structure in which Fe carbide ε is finely dispersed.
(2)前記Fe系炭化物εのサイズが1〜20nm、Fe系炭化物εの面積率が1.0%以上であることを特徴とする上記(1)に記載の耐遅れ破壊特性に優れた高強度鋼。 (2) The size of the Fe-based carbide ε is 1 to 20 nm and the area ratio of the Fe-based carbide ε is 1.0% or more. Strength steel.
(3)水素トラップ量が2.5ppm以上、限界拡散性水素量が6.0ppm以上であることを特徴とする上記(1)または(2)に記載の耐遅れ破壊特性に優れた高強度鋼。 (3) The high strength steel excellent in delayed fracture resistance according to the above (1) or (2), wherein the hydrogen trap amount is 2.5 ppm or more and the limit diffusible hydrogen amount is 6.0 ppm or more. .
(4)さらに、質量%で、
Cr:0.01〜5.00%、
Mo:0.01〜1.00%
のうちの1種もしくは2種を含有することを特徴とする上記(1)ないし(3)のいずれか1項に記載の耐遅れ破壊特性に優れた高強度鋼。
(4) Furthermore, in mass%,
Cr: 0.01 to 5.00%,
Mo: 0.01 to 1.00%
The high-strength steel excellent in delayed fracture resistance according to any one of the above (1) to (3), comprising one or two of the above.
(5)さらに、質量%で、
Nb:0.01〜0.10%、
V:0.01〜0.50%、
Ti:0.010〜0.300%、
Al:0.01〜0.20%
のうちの1種もしくは2種以上を含有することを特徴とする上記(1)ないし(4)のいずれか1項に記載の耐遅れ破壊特性に優れた高強度鋼。
(5) Furthermore, in mass%,
Nb: 0.01-0.10%,
V: 0.01 to 0.50%,
Ti: 0.010 to 0.300%,
Al: 0.01-0.20%
The high-strength steel excellent in delayed fracture resistance according to any one of the above (1) to (4), characterized by containing one or more of them.
(6)上記(1)ないし(5)のいずれに1項に記載の高強度鋼からなることを特徴とする高強度ボルト。 (6) A high-strength bolt comprising the high-strength steel according to any one of (1) to (5) above.
本発明によれば、上述した鋼成分範囲、熱処理条件、およびミクロ組織形態を選択することにより、耐遅れ破壊特性に優れた引張強さ1200MPaを超える高強度鋼を提供することが可能となる。このような高強度鋼は耐遅れ破壊特性の必要なボルトやばね等の各種構造材料として広範囲に適用が期待でき、産業上の効果は極めて顕著である。 According to the present invention, it is possible to provide a high-strength steel having a tensile strength exceeding 1200 MPa excellent in delayed fracture resistance by selecting the above-described steel component range, heat treatment conditions, and microstructure form. Such high-strength steel can be expected to be widely applied as various structural materials such as bolts and springs that require delayed fracture resistance, and the industrial effect is extremely remarkable.
本発明について、以下詳細に説明する。まず、上述した鋼成分範囲の限定理由について説明する。 The present invention will be described in detail below. First, the reason for limiting the steel component range described above will be described.
C:0.20〜1.50%
Cは鋼の強度を決める重要な元素である。十分に強度を得るためには、下限は0.20%とする。他の合金元素に比べて合金コストは安く、Cを多量に添加することができれば鋼材の合金コストは低減できる。しかしながら、多量のCを添加すると水素に対する脆化感受性が高まるため、上限は1.50%とする。
C: 0.20 to 1.50%
C is an important element that determines the strength of steel. In order to obtain sufficient strength, the lower limit is made 0.20%. Compared to other alloy elements, the alloy cost is low. If a large amount of C can be added, the alloy cost of the steel material can be reduced. However, the addition of a large amount of C increases the embrittlement susceptibility to hydrogen, so the upper limit is made 1.50%.
Si:0.20〜5.00%
SiはFe系炭化物εを安定化させ、Fe系炭化物θへの遷移を遅らせことで焼戻し軟化抵抗に大きく寄与するだけでなく、耐遅れ破壊特性を向上させる重要な元素である。これら効果を得るためには、下限を0.20%以上とする。好ましくは0.30%以上であり、より好ましくは0.50%以上である。しかしながら、多量に添加すると鋼材の表層に顕著な脱炭が生じるため、上限は5.00%とする。
Si: 0.20 to 5.00%
Si is an important element that not only greatly contributes to temper softening resistance by stabilizing the Fe-based carbide ε and delaying the transition to the Fe-based carbide θ, but also improves delayed fracture resistance. In order to obtain these effects, the lower limit is made 0.20% or more. Preferably it is 0.30% or more, more preferably 0.50% or more. However, if added in a large amount, remarkable decarburization occurs in the surface layer of the steel material, so the upper limit is made 5.00%.
Mn:0.10〜3.00%
Mnは鋼の焼入れ性を向上するのに有効な元素であるとともに、鋼中のSをMnSとして固定することによって熱間脆性を防止する効果がある。これら効果を得るためには、下限を0.10%以上とする。しかしながら、3.00%を超える添加は、かえって水素に対する脆化感受性が高め遅れ破壊特性を低下させる。したがって、上限を3.00%とする。
Mn: 0.10 to 3.00%
Mn is an element effective for improving the hardenability of steel and has an effect of preventing hot brittleness by fixing S in the steel as MnS. In order to obtain these effects, the lower limit is made 0.10% or more. However, the addition exceeding 3.00% increases the embrittlement susceptibility to hydrogen and lowers the delayed fracture characteristics. Therefore, the upper limit is made 3.00%.
P:0.0005〜0.1000%
Pは鋼中の不可避的不純物として、0.0005%以上は含有しているため、下限を0.0005%とする。鋼中のPは、耐遅れ破壊特性を低下させるため、上限を0.1000%とする。
P: 0.0005 to 0.1000%
Since P contains 0.0005% or more as an inevitable impurity in steel, the lower limit is made 0.0005%. P in steel lowers the delayed fracture resistance, so the upper limit is made 0.1000%.
S:0.0005〜0.2000%
SはPと同様に鋼中の不可避的不純物として通常、0.0005%以上は含有し、鋼中に存在すると鋼を脆化させる。Sの場合、一部MnSとして固定することによって、極力その影響は小さくなるものの、その添加量は少ないことが望ましい。したがって、Sの含有量は、0.0005〜0.2000%とする。
S: 0.0005 to 0.2000%
S, like P, usually contains 0.0005% or more as an unavoidable impurity in steel, and when present in steel, it causes embrittlement of the steel. In the case of S, it is desirable that the amount added is small although the influence is minimized by fixing as MnS partially. Therefore, the content of S is set to 0.0005 to 0.2000%.
N:0.0020〜0.0200%
Nは窒化物を形成し、旧オーステナイト粒を微細化する効果がある。この効果を得るためには、Nの下限は0.0020%とする。しかしながら、多量に添加しても、この効果は飽和するため、Nの上限は0.0200%とする。
N: 0.0020 to 0.0200%
N forms nitrides and has the effect of refining prior austenite grains. In order to obtain this effect, the lower limit of N is set to 0.0020%. However, even if added in a large amount, this effect is saturated, so the upper limit of N is 0.0200%.
Cr:0.01〜5.00%
Crは鋼の焼入れ性を向上するのに有効な元素であるとともに、焼戻し軟化抵抗に寄与する元素である。これら効果を得るためには、下限は0.01%とする。しかしながら、CrはFe系炭化物θ中に固溶し安定化させ、多量に添加した場合、焼入れ性の向上や焼戻し軟化抵抗の効果を得るために、焼入れ処理前に高温で加熱する必要がある。これにより、旧オーステナイト粒が粗大化し、耐遅れ破壊特性が低下する。したがって、上限は5.00%とする。
Cr: 0.01 to 5.00%
Cr is an element effective for improving the hardenability of steel and an element contributing to temper softening resistance. In order to obtain these effects, the lower limit is made 0.01%. However, Cr needs to be heated at a high temperature before the quenching treatment in order to obtain a solid solution and stabilize in the Fe-based carbide θ and to stabilize it and to add a large amount to obtain effects of improving hardenability and temper softening resistance. As a result, the prior austenite grains become coarse, and the delayed fracture resistance deteriorates. Therefore, the upper limit is made 5.00%.
Mo:0.01〜1.00%
Moは鋼の焼入れ性を向上するのに有効な元素であるとともに、焼戻し軟化抵抗に寄与する元素である。これら効果を得るためには、下限を0.01%とする。しかしながら、1.00%を超えて添加すると、これら効果は飽和する。したがって、上限を1.00%とする。
Mo: 0.01 to 1.00%
Mo is an element effective for improving the hardenability of steel and an element contributing to temper softening resistance. In order to obtain these effects, the lower limit is made 0.01%. However, these effects are saturated when added over 1.00%. Therefore, the upper limit is made 1.00%.
Nb:0.01〜0.10%、V:0.01〜0.50%、Ti:0.010〜0.300%、Al:0.01〜0.20%のうちの1種もしくは2種以上を含有する。
Nb、V、Ti、Alは窒化物を形成し、旧オーステナイト粒を微細化する効果がある。この効果を得るためには、Nb、V、Ti、Alの下限は0.01%とする。しかしながら、それぞれ多量に添加しても、この効果は飽和するため、Nbの上限は0.10%、Vの上限は0.50%、Tiの上限は0.300%、Alの上限は0.20%とする。
One or two of Nb: 0.01 to 0.10%, V: 0.01 to 0.50%, Ti: 0.010 to 0.300%, Al: 0.01 to 0.20% Contains more than seeds.
Nb, V, Ti, and Al have the effect of forming nitrides and refining prior austenite grains. In order to obtain this effect, the lower limit of Nb, V, Ti, and Al is 0.01%. However, even if added in a large amount, this effect is saturated. Therefore, the upper limit of Nb is 0.10%, the upper limit of V is 0.50%, the upper limit of Ti is 0.300%, and the upper limit of Al is 0.00%. 20%.
次に上述した鋼組織の限定理由について説明する。 Next, the reason for limiting the steel structure described above will be described.
鋼組織中の析出物をFe系炭化物εに規定したのは、粗大なFe系炭化物θでは外部環境から侵入した鋼中の水素をトラップすることができないのに対して、微細なFe系炭化物εは鋼中の水素をトラップすることができ、臨界拡散性水素量を増加させることができるためである。また鋼組織について焼戻しマルテンサイト組織に規定したのは、多量のFe系炭化物εを微細分散するためには、焼入れ焼戻し処理が有効であるためである。 The precipitates in the steel structure are defined as Fe-based carbides ε, because coarse Fe-based carbides θ cannot trap hydrogen in steel that has entered from the external environment, whereas fine Fe-based carbides ε. This is because hydrogen in steel can be trapped and the amount of critical diffusible hydrogen can be increased. The reason why the steel structure is defined as the tempered martensite structure is that quenching and tempering treatment is effective for finely dispersing a large amount of Fe-based carbide ε.
次に上述したFe系炭化物εのサイズや面積率の限定理由について説明する。 Next, the reasons for limiting the size and area ratio of the Fe-based carbide ε will be described.
Fe系炭化物εのサイズを1〜20nmに規定したのは、20nm超では母相と析出物の界面の整合性が保たれなくなり、水素をトラップする歪みが形成しないためであり、水素トラップ効果を得るには1nm以上が必要なためである。好ましくは2nm以上10nm以下である。またFe系炭化物εの面積率を1.0%以上に規定したのは、Fe系炭化物εが析出し水素をトラップしても、その量が少なければ耐遅れ破壊特性の大きな向上は得られないためである。上記面積率は、鋼中C量により自ずと飽和する。 The reason why the size of the Fe-based carbide ε is defined to be 1 to 20 nm is that when the thickness exceeds 20 nm, the consistency between the interface of the parent phase and the precipitate is not maintained, and the strain for trapping hydrogen is not formed. This is because a thickness of 1 nm or more is necessary to obtain. Preferably they are 2 nm or more and 10 nm or less. Also, the area ratio of Fe-based carbide ε is defined to be 1.0% or more. Even if Fe-based carbide ε precipitates and traps hydrogen, if the amount is small, the delayed fracture resistance cannot be greatly improved. Because. The area ratio is naturally saturated by the amount of C in steel.
次に上述した水素トラップ量と限界拡散性水素量の限定理由について説明する。 Next, the reasons for limiting the hydrogen trap amount and the limit diffusible hydrogen amount will be described.
水素トラップ量が2.5ppm未満、または限界拡散性水素量が6.0ppm未満では、外部環境から鋼材中に侵入する水素によって遅れ破壊が発生するため、水素トラップ量を2.5ppm以上、限界拡散性水素量を6.0ppm以上に限定した。 If the hydrogen trap amount is less than 2.5 ppm, or the limit diffusible hydrogen amount is less than 6.0 ppm, delayed fracture occurs due to hydrogen entering the steel from the external environment. The amount of reactive hydrogen was limited to 6.0 ppm or more.
次に上述した焼戻し温度の限定理由について説明する。 Next, the reason for limiting the tempering temperature described above will be described.
鋼に所定の強度および靱性、延性を付与するために、焼入れ後に焼戻しを行う必要がある。焼戻しは、一般に150℃からAC1点の温度範囲で行われるが、本発明では250〜525℃の温度範囲に限定する必要がある。その理由は、焼戻し温度が525℃を超えると、Siを5%添加しFe系炭化物εを安定化させても、Fe系炭化物θに遷移し臨界拡散性水素量が低下するためである。ただし、Fe系炭化物εからθに遷移する温度はSi量に大きく依存するため、規定した焼戻し温度範囲であったとしても、Si量が少なければ高温焼戻しでFe系炭化物θが析出する。 In order to impart predetermined strength, toughness and ductility to the steel, it is necessary to perform tempering after quenching. Tempering is generally performed in a temperature range from 150 ° C. to AC 1 point, but in the present invention, it is necessary to limit the temperature range to 250 to 525 ° C. The reason is that when the tempering temperature exceeds 525 ° C., even if 5% of Si is added to stabilize the Fe-based carbide ε, the transition to Fe-based carbide θ occurs and the critical diffusible hydrogen amount decreases. However, since the temperature at which the Fe-based carbide ε transitions to θ greatly depends on the amount of Si, even if the temperature is within the specified tempering temperature range, the Fe-based carbide θ is precipitated by high-temperature tempering if the Si amount is small.
そこでFe系炭化物εが析出する焼戻し処理条件とSi量の関係について検討した。C量:0.60%(質量%、以下同じ)、Si量:0.20〜5.00%、Mn量:0.75%、P量:0.005〜0.009%、S量:0.006〜0.0010%、N量:0.0070%、残部がFe及び不可避的不純物からなる鋼材を準備し、これら鋼材を1050℃に加熱後、焼戻し温度150〜550℃、保定時間1〜18000sec.の範囲で熱処理した。透過型電子顕微鏡を用いて鋼中の析出物を観察し、制限視野回折図形からFe系炭化物を特定した。検討の結果、Fe系炭化物εが析出する焼戻し処理条件と鋼中Si量との関係式(a)、(b)が得られた。
0.20≦[Si]<2.00 のとき
(273+T)×(log(t/3600)+20)
−2067[Si]−9467>0・・・(a)
2.00≦[Si]≦5.00 のとき
(273+T)×(log(t/3600)+20)
−733[Si]−12133>0・・・(b)
ここで、[Si]は鋼中のSi含有量(質量%)、Tは焼戻し温度(℃)、tは保定時間(秒)である。したがって、(a)式、及び(b)式の値が0以下の場合、Fe系炭化物θが析出する。また焼戻し温度525℃を超えると、Si量に関係なくFe系炭化物θが析出するため、(a)式、及び(b)式の適用範囲は焼戻し温度150〜525℃である。
Therefore, the relationship between the tempering conditions for precipitation of Fe-based carbide ε and the amount of Si was examined. C amount: 0.60% (mass%, the same applies hereinafter), Si amount: 0.20 to 5.00%, Mn amount: 0.75%, P amount: 0.005 to 0.009%, S amount: 0.006 to 0.0010%, N amount: 0.0070%, the balance is prepared with a steel material composed of Fe and inevitable impurities, and after heating these steel materials to 1050 ° C, a tempering temperature of 150 to 550 ° C and a holding time of 1 ~ 18000 sec. It heat-processed in the range. Precipitates in the steel were observed using a transmission electron microscope, and Fe-based carbides were identified from the limited field diffraction pattern. As a result of the study, relational expressions (a) and (b) between the tempering treatment conditions in which Fe-based carbide ε precipitates and the amount of Si in the steel were obtained.
When 0.20 ≦ [Si] <2.00 (273 + T) × (log (t / 3600) +20)
-2067 [Si] -9467> 0 (a)
When 2.00 ≦ [Si] ≦ 5.00 (273 + T) × (log (t / 3600) +20)
-733 [Si] -12133> 0 (b)
Here, [Si] is the Si content (% by mass) in the steel, T is the tempering temperature (° C.), and t is the retention time (seconds). Therefore, when the values of the formulas (a) and (b) are 0 or less, the Fe-based carbide θ is precipitated. Further, when the tempering temperature exceeds 525 ° C., Fe-based carbide θ precipitates regardless of the amount of Si, and therefore the applicable range of the formulas (a) and (b) is the tempering temperature 150 to 525 ° C.
ただし、焼戻し温度250℃未満では、Fe系炭化物εは析出するものの、母相中の転位等の欠陥密度が高く、水素に対する脆化感受性が高いため、本発明では、焼戻し下限温度を上記のとおり250℃とした。鋼中にFe系炭化物εが安定に存在するならば、より水素に対する脆化感受性を低下させるために、高温で焼戻しすることが望ましい。好ましくは、焼戻し温度350℃以上である。 However, when the tempering temperature is less than 250 ° C., the Fe-based carbide ε is precipitated, but the defect density such as dislocations in the matrix phase is high and the embrittlement sensitivity to hydrogen is high. Therefore, in the present invention, the lower tempering temperature is as described above. The temperature was 250 ° C. If the Fe-based carbide ε is stably present in the steel, it is desirable to temper at a high temperature in order to further reduce the susceptibility to hydrogen embrittlement. Preferably, the tempering temperature is 350 ° C. or higher.
本発明を実施例によって以下に詳述する。なお、これら実施例は本発明の技術的意義、効果を説明するためのものであり、本発明の範囲を限定するものではない。 The invention is described in detail below by means of examples. These examples are for explaining the technical significance and effects of the present invention, and do not limit the scope of the present invention.
表1に示す化学成分の鋼を真空溶解炉で溶製後、熱間圧延することによって15mmφの鋼棒材を作製した。その後、鋼成分組成に応じて、Fe系炭化物の溶体化や組織のオーステナイト化が可能な温度を850〜1050℃から選択して加熱し、60℃の油に焼入れした後、表2に示す焼戻し温度で60min、焼戻し処理を施した。これら熱処理材からJIS Z 2201の14号引張試験片を採取し、引張強さを評価した。 A steel bar having a diameter of 15 mm was produced by hot rolling the steel having chemical components shown in Table 1 in a vacuum melting furnace. Then, according to the steel component composition, the temperature at which solution of Fe-based carbides or austenitization of the structure can be selected from 850 to 1050 ° C. is heated and quenched into oil at 60 ° C. Tempering was performed for 60 minutes at a temperature. JIS Z 2201 No. 14 tensile test specimens were collected from these heat treated materials and evaluated for tensile strength.
鋼中の析出物を特定するために、透過型電子顕微鏡を用いて観察し、制限視野回折図形から判断した。析出物のサイズおよび面積率は、透過型電子顕微鏡で、1視野の面積を50000nm2とし、20視野観察し、画像解析にて平均値として求めたものである。 In order to identify precipitates in the steel, they were observed using a transmission electron microscope and judged from the limited-field diffraction pattern. The size and area ratio of the precipitates were determined as an average value by image analysis by observing 20 visual fields with a transmission electron microscope with an area of 1 visual field of 50000 nm 2 .
水素トラップ量は、3%チオシアン酸アンモニウム溶液にNaClを3g/l添加した水溶液に試験片を浸漬し0.2mA/cm2の電流密度で電解水素チャージを42時間行い、その後、室温で96時間放置した後、ガスクロマトグラフによる昇温水素分析法で測定した。ガスクロマトグラフの昇温速度は100℃/時間であり、室温から400℃までに試験片から放出される水素量を水素トラップ量と定義している。水素トラップ量を評価するための試験片は7mmφの丸棒を準備した。限界拡散性水素量は、鉄と鋼、Vol.83(1997)、p.454に記載の方法で測定した。限界拡散性水素量が6.0ppm未満であるものは遅れ破壊特性に劣ると判断した。 The amount of the hydrogen trap was determined by immersing the test piece in an aqueous solution obtained by adding 3 g / l of NaCl to a 3% ammonium thiocyanate solution and performing electrolytic hydrogen charging at a current density of 0.2 mA / cm 2 for 42 hours, and then at room temperature for 96 hours. After leaving it to stand, it was measured by a temperature rising hydrogen analysis method using a gas chromatograph. The temperature rising rate of the gas chromatograph is 100 ° C./hour, and the amount of hydrogen released from the test piece from room temperature to 400 ° C. is defined as the hydrogen trap amount. As a test piece for evaluating the amount of hydrogen trap, a 7 mmφ round bar was prepared. The amount of critical diffusible hydrogen is iron and steel, Vol. 83 (1997), p. Measured by the method described in 454. Those having a limit diffusible hydrogen content of less than 6.0 ppm were judged to be inferior in delayed fracture characteristics.
表2から分かるように、No.1〜16の本発明例は、いずれも鋼組成および鋼組織が規定範囲内であるため、引張強さ1600MPa以上の高強度にも関わらず遅れ破壊特性に優れる。また本発明例No.4と15、No.11と16を比較して分かるように、焼戻し温度が高いほど水素に対する脆化感受性が低下し、遅れ破壊特性に優れる。これに対して、比較例No.17及び20は、鋼組織は規定範囲内であり、水素をトラップする能力はあるものの、C、Mnの含有量が多く、水素に対する脆化感受性が高くなり、かえって遅れ破壊特性に劣る。また比較例No.18及び19は、Siの含有量が少なく、Fe系炭化物θが析出したため、遅れ破壊特性に劣る。また比較例No.21は、鋼組織は規定範囲内であり、水素をトラップする能力はあるものの、Crの含有量が多く、焼入れ性の向上や焼戻し軟化抵抗の効果を得るには、焼入れ処理前の加熱温度を1200℃の高温に加熱する必要がある。これにより、旧オーステナイト粒が粗大化し遅れ破壊特性に劣る。また比較例No.22は、焼戻し温度が550℃と高く、Fe系炭化物θが析出したため、遅れ破壊特性に劣る。また比較例No.23は、鋼組成および焼戻し温度が規定範囲内であるものの、Fe系炭化物εが析出する焼戻し条件とSi量の関係から逸脱したためにFe系炭化物θが析出し、遅れ破壊特性に劣る。 As can be seen from Table 2, no. In all of Examples 1 to 16 of the present invention, the steel composition and the steel structure are within the specified ranges, so that the delayed fracture property is excellent despite the high strength of 1600 MPa or more. In addition, Invention Example No. 4 and 15, no. As can be seen by comparing 11 and 16, the higher the tempering temperature, the lower the susceptibility to hydrogen embrittlement and the better the delayed fracture characteristics. In contrast, Comparative Example No. In Nos. 17 and 20, although the steel structure is within the specified range and has the ability to trap hydrogen, the contents of C and Mn are large, the embrittlement susceptibility to hydrogen is increased, and the delayed fracture characteristics are inferior. Comparative Example No. Nos. 18 and 19 are inferior in delayed fracture characteristics because the Si content is small and Fe-based carbides θ are precipitated. Comparative Example No. No. 21, although the steel structure is within the specified range and has the ability to trap hydrogen, the Cr content is large, and in order to obtain the effects of improved hardenability and temper softening resistance, the heating temperature before the quenching treatment should be adjusted. It is necessary to heat to a high temperature of 1200 ° C. Thereby, the prior austenite grains become coarse and inferior in delayed fracture characteristics. Comparative Example No. No. 22 has a high tempering temperature of 550 ° C., and the Fe-based carbide θ is precipitated. Comparative Example No. In No. 23, although the steel composition and the tempering temperature are within the specified ranges, the deviation from the relationship between the tempering conditions in which the Fe-based carbide ε precipitates and the amount of Si causes Fe-based carbide θ to precipitate, and the delayed fracture characteristics are inferior.
Claims (6)
C:0.20〜1.50%、
Si:0.20〜5.00%、
Mn:0.10〜3.00%、
P:0.0005〜0.1000%、
S:0.0005〜0.2000%、
N:0.0020〜0.0200%
を含有し、残部がFe及び不可避的不純物からなり、鋼組織はFe系炭化物εが微細分散した焼戻しマルテンサイト組織であることを特徴とする耐遅れ破壊特性に優れた高強度鋼。 % By mass
C: 0.20 to 1.50%,
Si: 0.20 to 5.00%,
Mn: 0.10 to 3.00%,
P: 0.0005 to 0.1000%,
S: 0.0005 to 0.2000%,
N: 0.0020 to 0.0200%
A high-strength steel excellent in delayed fracture resistance, characterized in that the balance is made of Fe and inevitable impurities, and the steel structure is a tempered martensite structure in which Fe carbide ε is finely dispersed.
Cr:0.01〜5.00%、
Mo:0.01〜1.00%
のうちの1種もしくは2種を含有することを特徴とする請求項1ないし3のいずれか1項に記載の耐遅れ破壊特性に優れた高強度鋼。 Furthermore, in mass%,
Cr: 0.01 to 5.00%,
Mo: 0.01 to 1.00%
The high-strength steel excellent in delayed fracture resistance according to any one of claims 1 to 3, wherein one or two of them are contained.
Nb:0.01〜0.10%、
V:0.01〜0.50%、
Ti:0.010〜0.300%、
Al:0.01〜0.20%、
のうちの1種もしくは2種以上を含有することを特徴とする請求項1ないし4のいずれか1項に記載の耐遅れ破壊特性に優れた高強度鋼。 Furthermore, in mass%,
Nb: 0.01-0.10%,
V: 0.01 to 0.50%,
Ti: 0.010 to 0.300%,
Al: 0.01-0.20%,
The high-strength steel excellent in delayed fracture resistance according to any one of claims 1 to 4, characterized by containing one or more of them.
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