JP2006045670A - High strength heat treated steel having excellent delayed fracture resistance and production method therefor - Google Patents
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 39
- 239000010959 steel Substances 0.000 title claims abstract description 39
- 230000003111 delayed effect Effects 0.000 title claims abstract description 34
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 10
- 229910045601 alloy Inorganic materials 0.000 claims abstract description 40
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- 229910052719 titanium Inorganic materials 0.000 claims abstract description 13
- 229910052758 niobium Inorganic materials 0.000 claims abstract description 12
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 9
- 229910052750 molybdenum Inorganic materials 0.000 claims abstract description 8
- 229910052720 vanadium Inorganic materials 0.000 claims abstract description 7
- 239000012535 impurity Substances 0.000 claims abstract description 4
- 229910052739 hydrogen Inorganic materials 0.000 claims description 48
- 239000001257 hydrogen Substances 0.000 claims description 48
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 claims description 47
- 238000005496 tempering Methods 0.000 claims description 19
- 239000000463 material Substances 0.000 claims description 9
- 238000010791 quenching Methods 0.000 claims description 9
- 230000000171 quenching effect Effects 0.000 claims description 9
- 229910052804 chromium Inorganic materials 0.000 claims description 2
- 229910052802 copper Inorganic materials 0.000 claims description 2
- 229910052748 manganese Inorganic materials 0.000 claims description 2
- 229910052759 nickel Inorganic materials 0.000 claims description 2
- 229910052757 nitrogen Inorganic materials 0.000 claims description 2
- 150000001247 metal acetylides Chemical class 0.000 abstract description 8
- 229910052799 carbon Inorganic materials 0.000 abstract description 3
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- 238000005728 strengthening Methods 0.000 description 10
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- 238000010438 heat treatment Methods 0.000 description 3
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- 230000002730 additional effect Effects 0.000 description 1
- SOIFLUNRINLCBN-UHFFFAOYSA-N ammonium thiocyanate Chemical compound [NH4+].[S-]C#N SOIFLUNRINLCBN-UHFFFAOYSA-N 0.000 description 1
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- 125000004429 atom Chemical group 0.000 description 1
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- 125000004435 hydrogen atom Chemical group [H]* 0.000 description 1
- 239000004615 ingredient Substances 0.000 description 1
- 229910052742 iron Inorganic materials 0.000 description 1
- 239000002184 metal Substances 0.000 description 1
- 229910052751 metal Inorganic materials 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 229910001562 pearlite Inorganic materials 0.000 description 1
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Abstract
Description
本発明は、耐遅れ破壊特性に優れた高強度調質鋼およびその製造方法に関し、特に、ボルト用鋼を代表とする、引張り強さが１２００〜１６００ＭＰａの耐遅れ破壊特性に優れた高強度調質鋼およびその製造方法に関するものである。 The present invention relates to a high-strength tempered steel excellent in delayed fracture resistance and a method for producing the same, and in particular, a high-strength tempered steel excellent in delayed fracture resistance with a tensile strength of 1200 to 1600 MPa, represented by bolt steel. The present invention relates to a quality steel and a manufacturing method thereof.
鋼材の高強度化にともない、耐遅れ破壊に対する対策は産業利用上、安全性確保の立場から必須であり、これまでに特許文献１、特許文献２等に記載された先行技術に代表されるように、耐遅れ破壊特性の優れた高強度鋼が開発されてきた。 With increasing strength of steel materials, countermeasures against delayed fracture resistance are indispensable for industrial use from the standpoint of ensuring safety, and so far as represented by the prior art described in Patent Document 1, Patent Document 2, etc. In addition, high strength steels having excellent delayed fracture resistance have been developed.
耐遅れ破壊特性を向上させる技術として、代表的なものの一つに、調質処理において、焼入後に合金炭化物の安定領域のなかでも高温域で焼戻しを行なうことが挙げられる。高温域で焼戻しすることにより、粒界にフィルム状に存在する炭化物を球状化させ、粒界破壊を阻止する、また、マルテンサイト中に合金炭化物を微細析出させることにより、二次硬化が発生すると共に、合金炭化物の水素トラップ効果で、鋼中水素の拡散・集積を妨げることにより遅れ破壊を起こし難くすることが知られている。 One typical technique for improving delayed fracture resistance is to perform tempering in the high temperature region of the alloy carbide in the stable region after quenching in the tempering treatment. By tempering in a high temperature region, the carbide present in the form of a film at the grain boundary is spheroidized to prevent grain boundary destruction, and secondary hardening occurs by alloy fine carbide precipitated in martensite. At the same time, it is known that delayed fracture is less likely to occur by hindering the diffusion and accumulation of hydrogen in steel due to the hydrogen trap effect of alloy carbides.
しかし、従来技術の問題点として次のような点が挙げられる。まず第１に、既存の耐遅れ破壊特性に優れた高強度ボルト用鋼は上記のような考え方から、合金炭化物を形成する元素を多量に添加することが行われてきたが、必ずしも炭素量とのバランス上、最適な成分となっていなかった。 However, there are the following points as problems of the prior art. First of all, existing steels for high-strength bolts with excellent delayed fracture resistance have been added in a large amount of elements that form alloy carbides based on the above concept. In terms of balance, it was not an optimal ingredient.
第２に、合金炭化物を水素トラップとして有効に作用させるには、マルテンサイト中に合金炭化物を微細析出されることが必要であり、そのために製造工程において一旦溶体化処理を行なった後、焼き入れ、焼戻しを行なうことが必要であるが、必ずしも炭化物の熱安定性を評価して、熱処理条件を決定していなかった。 Secondly, in order to make the alloy carbide effectively act as a hydrogen trap, it is necessary to finely precipitate the alloy carbide in the martensite. Therefore, after the solution treatment is once performed in the manufacturing process, quenching is performed. Although tempering is necessary, the thermal stability of the carbide has not necessarily been evaluated and the heat treatment conditions have not been determined.
本発明は、上述した課題を解決し、従来と比較して水素トラップおよび析出強化量を効果的に発現するための最適な炭化物のサイズ、体積率、整合歪を見出し、それを得るための合金添加量および熱処理条件を示し、耐遅れ破壊特性に優れた高強度調質鋼およびその製造方法を提供するものである。 The present invention solves the above-mentioned problems, finds the optimum carbide size, volume ratio, and matching strain for effectively expressing the hydrogen trap and precipitation strengthening amount as compared with the conventional alloy, and an alloy for obtaining the same The present invention provides a high-strength tempered steel excellent in delayed fracture resistance, and a method for producing the same, showing the addition amount and heat treatment conditions.
本発明は上記の課題を解決するためになされたものであり、その趣旨とするところは次の通りである。
（１）質量％で、Ｃ：０．１〜０．５％を含有し、更に、Ｖ≧０．０５％、Ｍｏ：０．１％以上、３．０％未満、Ｔｉ≧０．０３％、Ｎｂ≧０．０５％のうち１種または２種以上を含有し、且つＣとの質量％比が、０．５≦（０．１８Ｖ＋０．０６Ｍｏ＋０．２５Ｔｉ＋０．１３Ｎｂ）／Ｃであり、残部がＦｅおよび不可避不純物からなり、引張強さが１２００〜１６００ＭＰａ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調質鋼。
（２）さらに、質量％で、Ｓｉ：０．０５〜２％、Ｍｎ：０．０５〜３％、Ｎ≦０．００８％、Ａｌ：０．００５〜０．１％、Ｃｕ≦０．３％、Ｎｉ≦２％、Ｃｒ≦２．５％、Ｂ≦０．０１％の１種または２種以上を含有す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調質鋼。
（３）焼戻しマルテンサイト中の合金炭化物サイズが５〜４０ｎｍ、合金炭化物の体積率が０．３％以上、合金炭化物の整合歪が１１％以下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調質鋼。
（４）水素トラップ量が２ｐｐｍ以上、限界拡散性水素量が５ｐｐｍ以上であ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調質鋼。
（５）上記（１）または（２）に記載の鋼成分を有する鋼材を、鋼材の成分としてＴｉまたはＮｂのいずれか１種類または両方を含有するものについては１１００℃以上の温度にて溶体化処理後に焼入れし、Ｔｉ、Ｎｂのいずれも含有しないものについては９５０℃以上の温度にて溶体化処理後に焼入れし、その後５５０〜７００℃の温度範囲で焼戻すことを特徴とする、耐遅れ破壊特性に優れた高強度調質鋼の製造方法。
The present invention has been made to solve the above problems, and the gist of the present invention is as follows.
(1) By mass%, containing C: 0.1-0.5%, V ≧ 0.05%, Mo: 0.1% or more, less than 3.0%, Ti ≧ 0.03% , Nb ≧ 0.05%, and the mass% ratio with C is 0.5 ≦ (0.18V + 0.06Mo + 0.25Ti + 0.13Nb) / C, and the balance is A high-strength tempered steel excellent in delayed fracture resistance, characterized by comprising Fe and inevitable impurities and having a tensile strength of 1200 to 1600 MPa.
(2) Further, by mass, Si: 0.05-2%, Mn: 0.05-3%, N ≦ 0.008%, Al: 0.005-0.1%, Cu ≦ 0.3 %, Ni ≦ 2%, Cr ≦ 2.5%, B ≦ 0.01%, one or two or more types, characterized by excellent delayed fracture resistance according to (1) above High strength tempered steel.
(3) The alloy carbide size in the tempered martensite is 5 to 40 nm, the volume fraction of the alloy carbide is 0.3% or more, and the matching strain of the alloy carbide is 11% or less, (1) or High strength tempered steel with excellent delayed fracture resistance described in (2).
(4) The hydrogen trap amount is 2 ppm or more and the critical diffusible hydrogen amount is 5 ppm or more, and the high resistance to delayed fracture resistance according to any one of (1) to (3) above Strength tempered steel.
(5) A steel material having the steel component described in (1) or (2) above is solution-treated at a temperature of 1100 ° C. or higher for a steel material containing either one or both of Ti and Nb. Quenching after the treatment, for those containing neither Ti nor Nb, delayed hardening resistance characterized by quenching after solution treatment at a temperature of 950 ° C. or higher and then tempering in a temperature range of 550 to 700 ° C. A method for producing high strength tempered steel with excellent properties.
本発明は、請求項に示す成分範囲の鋼において、所定の製造方法に従って処理することにより、上記の金属組織上の特徴を有し、１２００〜１６００ＭＰａの高強度が得られるとともに、耐遅れ破壊特性にも優れた鋼材が実現した。このことで、高強度調質鋼の耐遅れ破壊特性を大幅に向上させたものであり、産業上の効果は極めて顕著である。 The present invention is characterized in that the steel having the component ranges shown in the claims is processed according to a predetermined manufacturing method and has the above-mentioned metal structure characteristics, and high strength of 1200 to 1600 MPa is obtained, and delayed fracture resistance Excellent steel material was realized. As a result, the delayed fracture resistance of the high-strength tempered steel is greatly improved, and the industrial effect is extremely remarkable.
本発明において、化学組成範囲の意味について述べる。 In the present invention, the meaning of the chemical composition range will be described.
Ｃは、鋼の強度を向上させる有効な成分として添加するもので、０．１％未満では調質しても、遅れ破壊が問題となる強度域である１２００ＭＰａ以上までに到達しないため下限を０．１％に限定した。また、０．５％超の鋼の調質では、強度が１６００ＭＰａを越え、遅れ破壊以前の問題として破断面形態がへき開へと移行し、十分な靭性が得られないため、上限を０．５％に限定した。 C is added as an effective component for improving the strength of the steel, and even if it is tempered at less than 0.1%, the lower limit is 0 because it does not reach 1200 MPa or more, which is a strength region where delayed fracture is a problem. Limited to 1%. In addition, when the tempering of steel exceeds 0.5%, the strength exceeds 1600 MPa, and the fracture surface shape shifts to cleavage as a problem before delayed fracture, and sufficient toughness cannot be obtained. %.
Ｖ、Ｍｏ、Ｔｉ、Ｎｂは、いずれも焼戻しの過程において、マルテンサイト組織中に合金炭化物として微細析出し、高強度化および耐遅れ破壊特性の向上に寄与する効果を有している。Ｖ：０．０５％未満、Ｍｏ：０．１％未満、Ｔｉ：０．０３％未満、Ｎｂ：０．０５％未満では、上記特性の向上効果が十分に発揮できず、他方、Ｍｏは高価な元素であり、多量の添加は経済性の観点から好ましくないため、Ｖ≧０．０５％、Ｍｏ：０．１％以上、３．０％未満、Ｔｉ≧０．０３％、Ｎｂ≧０．０５％に限定した。 V, Mo, Ti, and Nb are all finely precipitated as alloy carbides in the martensite structure during the tempering process, and have the effect of contributing to high strength and improved delayed fracture resistance. When V is less than 0.05%, Mo is less than 0.1%, Ti is less than 0.03%, and Nb is less than 0.05%, the effect of improving the above characteristics cannot be sufficiently exhibited. On the other hand, Mo is expensive. Since addition of a large amount is not preferable from the viewpoint of economy, V ≧ 0.05%, Mo: 0.1% or more, less than 3.0%, Ti ≧ 0.03%, Nb ≧ 0. Limited to 05%.
Ｖ、Ｍｏ、Ｔｉ、Ｎｂの添加量が、合金炭化物を生成する際に化学量論的にＣ量に対して不足する場合、マトリックスに残存固溶するＣが焼戻し時に水素のトラップサイトである炭化物の整合析出物／マトリックス界面、および界面近傍に分布するマトリックス中の整合歪場に偏析し、水素のトラップサイト能を低減させる。また、同時にこの条件では、析出強化に有効な合金炭化物の生成が不充分であり、高強度化に対する寄与も不充分となることから合金元素を質量％比で０．５≦（０．１８Ｖ＋０．０６Ｍｏ＋０．２５Ｔｉ＋０．１３Ｎｂ）／Ｃに限定した。また、合金元素の添加量をさらに増加することによって合金炭化物の生成がより促進されるという理由から、より好ましくは０．６≦（０．１８Ｖ＋０．０６Ｍｏ＋０．２５Ｔｉ＋０．１３Ｎｂ）／Ｃである。 When the addition amount of V, Mo, Ti, and Nb is stoichiometrically insufficient with respect to the amount of C when forming alloy carbide, the carbide that remains as a solid solution in the matrix is a hydrogen trap site during tempering. Segregated in the matching strain field in the matrix of the coherent precipitate / matrix and in the matrix distributed in the vicinity of the interface, thereby reducing the hydrogen trap site ability. At the same time, under these conditions, the formation of alloy carbides effective for precipitation strengthening is insufficient, and the contribution to increasing the strength is insufficient, so that the alloy elements are contained in a mass% ratio of 0.5 ≦ (0.18V + 0. 06Mo + 0.25Ti + 0.13Nb) / C. Further, it is more preferable that 0.6 ≦ (0.18V + 0.06Mo + 0.25Ti + 0.13Nb) / C because the formation of alloy carbide is further promoted by further increasing the amount of alloy element added.
逆に、Ｖ、Ｍｏ、Ｔｉ、ＮｂがＣに対して化学量論組成より一定量以上過剰に合金元素を添加しても、析出強化量、水素トラップ量の向上効果が飽和するため、合金添加量の質量％比である（０．１８Ｖ＋０．０６Ｍｏ＋０．２５Ｔｉ＋０．１３Ｎｂ）／Ｃが１．５を越えない範囲とするのが好ましい。 On the contrary, V, Mo, Ti, Nb is added to the alloy because the effect of improving the precipitation strengthening amount and the amount of hydrogen trap is saturated even if the alloy element is added in excess of a certain amount over the stoichiometric composition with respect to C. It is preferable that the mass% ratio (0.18V + 0.06Mo + 0.25Ti + 0.13Nb) / C not exceed 1.5.
Ｓｉは、固溶体硬化作用によって引張強さを高める作用があり、必要に応じて添加する。０．０５％未満では前記作用が発揮できず、一方、２％を超えても添加量に見合う効果が期待できないため、０．０５〜２％の範囲に限定した。 Si has the effect of increasing the tensile strength by the solid solution hardening effect, and is added as necessary. If the content is less than 0.05%, the above-described effect cannot be exhibited. On the other hand, if the content exceeds 2%, an effect commensurate with the amount of addition cannot be expected.
Ｍｎは、脱酸、脱硫のために、また、焼入性を向上させるのに有効な元素であり、必要に応じて添加するが、０．０５％未満では上記の効果が得られず、一方３％を越えて添加しても効果が飽和するため、０．０５〜３％に限定した。 Mn is an element effective for deoxidation and desulfurization and for improving hardenability, and is added as necessary. However, if it is less than 0.05%, the above effect cannot be obtained. Even if added over 3%, the effect is saturated, so it was limited to 0.05 to 3%.
Ｎは、製鋼工程で不可避的に不純物として含有されるが、０．００８％を越えて含まれると、Ｖ、Ｎｂ、Ｔｉの窒化物を生成し、微細炭化物生成による当該効果の発現を阻害するため、上限を０．００８％に規定した。 N is inevitably contained as an impurity in the steelmaking process, but if it exceeds 0.008%, nitrides of V, Nb, and Ti are generated, and the expression of the effect due to fine carbide generation is inhibited. Therefore, the upper limit is defined as 0.008%.
Ａｌは、Ｎを固定する効果があり、必要に応じて添加するが、０．００５％未満では殆どその効果を発現せず、０．１％を越えて添加しても添加量に見合う効果が期待出来ないため、０．００５〜０．１％に限定した。 Al has an effect of fixing N, and is added as necessary. However, if it is less than 0.005%, the effect is hardly expressed, and even if added over 0.1%, there is an effect commensurate with the addition amount. Since it cannot be expected, the content is limited to 0.005 to 0.1%.
Ｃｕは、焼き戻しの過程で炭化物とは独立に析出し、析出強化の追加効果を示す元素であり、必要に応じて添加するが、０．００５％未満ではその効果が薄いため、０．００５％以上とするのが望ましい。また、Ｃｕは、熱間で鋼材を加工する際に表面に割れ疵を生じさせる元素であり、０．３％を越えて添加すると割れ疵の発生する頻度が著しく増大する。このため、上限を０．３％に限定した。 Cu is an element that precipitates independently of the carbide during the tempering process and exhibits an additional effect of precipitation strengthening, and is added as necessary. However, if less than 0.005%, the effect is small, so 0.005 % Or more is desirable. Further, Cu is an element that generates cracks on the surface when the steel material is processed hot, and if added over 0.3%, the frequency of occurrence of cracks increases remarkably. For this reason, the upper limit was limited to 0.3%.
Ｎｉは、素地の低温靱性を高める効果があり、必要に応じて添加するが、０．００５％未満ではその効果を発揮しないため、０．００５％以上とするのが望ましい。また、Ｎｉは、冷間での加工性を阻害する傾向があり、２％を越えて添加すると冷間鍛造性、および切削性を著しく悪化させるため、上限を２％に規定した。 Ni has the effect of increasing the low-temperature toughness of the substrate, and is added as necessary. However, if it is less than 0.005%, the effect is not exhibited, so 0.005% or more is desirable. Ni has a tendency to inhibit the workability in the cold, and if added over 2%, the cold forgeability and the machinability are remarkably deteriorated, so the upper limit is defined as 2%.
Ｃｒは、少量の添加により高温焼戻しで軟化の遅滞を生じさせる効果があり、必要に応じて添加するが、０．００５％未満ではその効果を発揮しないため、０．００５％以上とするのが望ましい。また、Ｃｒは、Ｍｏ2Ｃ、ＶＣの微細な炭化物を消失させ、むしろ軟化させる傾向があり、２．５％を越えて添加すると、その弊害が現れるため、上限を２．５％に規定した。 Cr has the effect of causing a delay in softening by high-temperature tempering when added in a small amount, and is added as necessary. However, if less than 0.005%, the effect is not exerted, so 0.005% or more is recommended. desirable. Further, Cr tends to cause the fine carbides of Mo 2 C and VC to disappear and rather soften, and if added over 2.5%, its adverse effect appears, so the upper limit was specified to 2.5%. .
Ｂは、焼入性を高めるのに有効な元素であり、必要に応じて添加するが、０．０００１％未満ではその効果を発揮しないため、０．０００１％以上とするのが望ましい。また、Ｂは、熱間での加工時にぜい性を発現させる元素である。０．０１％を越えて添加すると、その弊害が現れるため、上限を０．０１％に規定した。 B is an element effective for improving hardenability, and is added as necessary. However, if it is less than 0.0001%, the effect is not exhibited, so 0.0001% or more is desirable. B is an element that develops brittleness during hot working. If added over 0.01%, its adverse effect appears, so the upper limit was defined as 0.01%.
Ｐ、Ｓは、高強度鋼の耐遅れ破壊特性を向上させる観点から、それぞれ０．０１５％以下が好ましい範囲である。 P and S are each preferably in a range of 0.015% or less from the viewpoint of improving delayed fracture resistance of high-strength steel.
次に合金炭化物のサイズを規定した理由を述べる。本発明鋼において析出強化が極大となる析出物の直径は５〜４０ｎｍである。５ｎｍに満たない合金炭化物、および４０ｎｍを超える合金炭化物の場合、水素トラップおよび析出強化が不充分となることから、合金炭化物のサイズを５〜４０ｎｍに限定した。 Next, the reason for defining the size of the alloy carbide will be described. The diameter of the precipitate where precipitation strengthening is maximized in the steel of the present invention is 5 to 40 nm. In the case of alloy carbides less than 5 nm and alloy carbides exceeding 40 nm, the size of the alloy carbide was limited to 5 to 40 nm because hydrogen trapping and precipitation strengthening were insufficient.
次に合金炭化物の体積率を規定した理由を述べる。析出物サイズが上述の５〜４０ｎｍの範囲内で分布していれば、体積率の増加に伴って析出強化量は単調に増加するが、合金炭化物の体積率が０．３％未満である場合は、析出強化および水素トラップへの寄与が不充分であるため、下限値を０．３％に限定した。 Next, the reason for defining the volume fraction of alloy carbide will be described. If the precipitate size is distributed within the range of 5 to 40 nm as described above, the precipitation strengthening amount increases monotonously as the volume ratio increases, but the volume ratio of the alloy carbide is less than 0.3%. Has insufficient contribution to precipitation strengthening and hydrogen trapping, so the lower limit was limited to 0.3%.
次に、合金炭化物／マトリックス間の整合歪を限定した理由を述べる。転位はその線張力で析出物およびその周辺の応力場を切断して進むので、整合歪が大きいほど高強度化への寄与が増加することに加えて、整合歪により形成される応力場が水素トラップサイトとして機能する。整合歪が１１％を超える場合、整合条件が崩れ一部が非整合となる界面が存在するため、析出強化にはほとんど寄与せず、水素トラップ量が低下する。従って、整合歪を１１％以下に限定した。 Next, the reason for limiting the matching strain between the alloy carbide / matrix will be described. Since dislocation proceeds by cutting the precipitate and its surrounding stress field with its linear tension, the greater the matching strain, the greater the contribution to higher strength. Functions as a trap site. When the alignment strain exceeds 11%, the alignment condition is broken and there is an interface where a part of the alignment strain becomes non-aligned. Therefore, the alignment strain hardly contributes to precipitation strengthening, and the amount of hydrogen traps decreases. Therefore, the matching strain is limited to 11% or less.
本発明で限定している合金炭化物のサイズと体積率は、電界放射型透過電子顕微鏡で１０視野以上を観察し、平均値として求めたものである。また、整合歪は電界放射型電子顕微鏡で合金炭化物／マトリックス界面の格子像を撮影し、界面とマトリックスの格子定数を測定し求めたものである。 The size and volume ratio of the alloy carbide limited in the present invention are obtained as an average value by observing 10 or more fields with a field emission transmission electron microscope. The matching strain is obtained by taking a lattice image of the alloy carbide / matrix interface with a field emission electron microscope and measuring the lattice constant of the interface and the matrix.
次に水素トラップ量と限界拡散性水素量については、水素トラップ量が２ｐｐｍ未満または限界拡散性水素量が５ｐｐｍ未満では、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２ｐｐｍ以上、限界拡散性水素量を５ｐｐｍ以上に限定した。 Next, with regard to the hydrogen trap amount and the limit diffusible hydrogen amount, if the hydrogen trap amount is less than 2 ppm or the limit diffusible hydrogen amount is less than 5 ppm, delayed fracture occurs due to hydrogen entering the steel from the environment. The amount was limited to 2 ppm or more, and the limit diffusible hydrogen content was limited to 5 ppm or more.
次に、熱処理条件を限定した理由を述べる。Ｔｉ、Ｎｂのうちいずれか１種類または両方を含有するものについては、溶体化温度が１１００℃未満では、粗大な合金炭化物が未固溶で残存し、マトリックス中に固溶するＣ量、及び合金元素量が少なくなるので、その後の焼戻し時の微細な合金炭化物の析出量が少なくなり、また、マトリックス中に固溶するＣ量が低下することにより、焼入れ後のマルテンサイトが十分な焼き入れ硬さに到達しないため、下限を１１００℃に設定した。Ｔｉ、Ｎｂのいずれも含有しないものについても、同様に、溶体化温度が９５０℃未満では、粗大な合金炭化物が未固溶で残存し、マトリックス中に固溶するＣ量、及び合金元素量が少なくなるので、その後の焼戻し時の微細な合金炭化物の析出量が少なくなり、また、マトリックス中に固溶するＣ量が低下することにより、焼入れ後のマルテンサイトが十分な焼き入れ硬さに到達しないため、下限を９５０℃に設定した。なお、Ｔｉ、Ｎｂの含有に関わらず、工業的に短時間で溶体化を行うためには１２００℃以上で溶体化を行うのがより好ましい。ここで、Ｔｉを含有とはＴｉ≧０．０３％を意味し、Ｎｂを含有とはＮｂ≧０．０５％を意味する。 Next, the reason why the heat treatment conditions are limited will be described. For those containing either one or both of Ti and Nb, if the solution temperature is less than 1100 ° C., the coarse alloy carbide remains undissolved and the amount of C dissolved in the matrix, and the alloy Since the amount of element decreases, the amount of fine alloy carbide precipitated during subsequent tempering decreases, and the amount of C dissolved in the matrix decreases, so that the martensite after quenching is sufficiently hardened. Therefore, the lower limit was set to 1100 ° C. In the case where neither Ti nor Nb is contained, similarly, when the solution temperature is less than 950 ° C., the coarse alloy carbide remains undissolved, the amount of C dissolved in the matrix, and the amount of alloy elements Since the amount of precipitation of fine alloy carbide during subsequent tempering decreases, and the amount of C dissolved in the matrix decreases, martensite after quenching reaches sufficient quenching hardness. Therefore, the lower limit was set to 950 ° C. Regardless of the contents of Ti and Nb, it is more preferable to carry out the solution treatment at 1200 ° C. or higher in order to industrially carry out the solution solution in a short time. Here, containing Ti means Ti ≧ 0.03%, and containing Nb means Nb ≧ 0.05%.
焼戻し温度が５５０℃未満では原子の拡散が不十分であるために、合金炭化物が微細に析出することができないため下限を５５０℃に制限した。また焼戻し温度が７００℃を越えると、合金炭化物が成長して粗大化し、水素トラップ能や析出強化能を失ってしまうため、焼戻し温度の上限を７００℃に限定した。 When the tempering temperature is lower than 550 ° C., the diffusion of atoms is insufficient, and the alloy carbide cannot be finely precipitated, so the lower limit is limited to 550 ° C. Further, when the tempering temperature exceeds 700 ° C., the alloy carbide grows and becomes coarse and loses the hydrogen trapping ability and the precipitation strengthening ability. Therefore, the upper limit of the tempering temperature is limited to 700 ° C.
本発明の水素トラップ量は、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料を浸漬し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℃までに試料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。また、水素トラップ量を評価するための試験片は、図１に示す形状の丸棒試験片を用いた。 The amount of hydrogen trap of the present invention is such that the sample is immersed in an aqueous solution in which 3 g / l of NaCl is added to a 3% ammonium thiocyanate solution, and electrolytic hydrogen charging is performed at a current density of 0.2 mA / cm 2 for 42 hours, and then at room temperature. After being allowed to stand for 96 hours, 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 sample from room temperature to 400 ° C. is defined as the hydrogen trap amount. Moreover, the round bar test piece of the shape shown in FIG. 1 was used for the test piece for evaluating the amount of hydrogen traps.
限界拡散性水素量は、電解水素チャージにより種々のレベルの拡散性水素量を試料に含有させた後、遅れ破壊試験中に試料から大気中に水素が抜けることを防止するためにＣｄめっきを施し、その後、大気中で所定の荷重を負荷し、遅れ破壊が発生しなくなる上限の拡散性水素量を評価するものである。ここで、遅れ破壊試験片は図２に示す形状の切欠き形状のものであり、遅れ破壊試験の荷重は最大引張荷重の０．９倍である。遅れ破壊試験において、１００時間経過しても遅れ破壊しない最大の水素量を限界拡散性水素とした。昇温水素分析法によりで室温から４００℃までに放出される水素量を測定した。 The critical diffusible hydrogen amount is obtained by adding a diffusible hydrogen amount of various levels to the sample by electrolytic hydrogen charging, and then performing Cd plating to prevent hydrogen from escaping from the sample to the atmosphere during the delayed fracture test. Thereafter, the upper limit of the amount of diffusible hydrogen is evaluated by applying a predetermined load in the atmosphere so that delayed fracture does not occur. Here, the delayed fracture test piece has a notch shape shown in FIG. 2, and the load in the delayed fracture test is 0.9 times the maximum tensile load. In the delayed fracture test, the maximum hydrogen content that does not cause delayed fracture even after 100 hours has elapsed is defined as critical diffusible hydrogen. The amount of hydrogen released from room temperature to 400 ° C. was measured by a temperature rising hydrogen analysis method.
以下、実施例により本発明の効果をさらに具体的に説明する。表１に示す化学成分を有する供試材を通常の熱間圧延条件で圧延した。その後、溶体化処理を行い水冷または油冷により焼入れ処理を行い、表１に示す温度で焼戻し処理を行い、引張試験片、図１および図２に示す試験片に機械加工した。ミクロ組織は、いずれも焼戻しマルテンサイトが面積率で９６〜１００％であり、残部はフェライト、ベイナイト、パーライトの１種又は２種以上であった。 Hereinafter, the effects of the present invention will be described more specifically with reference to examples. Sample materials having chemical components shown in Table 1 were rolled under normal hot rolling conditions. Thereafter, a solution treatment was performed, a quenching treatment was performed by water cooling or oil cooling, a tempering treatment was performed at a temperature shown in Table 1, and a tensile test piece and a test piece shown in FIGS. 1 and 2 were machined. As for the microstructure, tempered martensite was 96-100% in area ratio in all cases, and the balance was one or more of ferrite, bainite, and pearlite.
上記の試料を用いて、機械的性質、水素トラップ量および限界拡散性水素量を評価した結果を表１に併せて示す。表１の試験Ｎｏ．１〜３２が本発明例で、試験Ｎｏ．３３〜４０が比較例である。同表に見られるように本発明例は、いずれも水素トラップ量が２ｐｐｍ以上であるとともに限界拡散性水素量が５ｐｐｍ以上であり、耐遅れ破壊特性の優れた１２００〜１６００ＭＰａの高強度鋼が実現されている。 Table 1 also shows the results of evaluating the mechanical properties, the amount of hydrogen trap and the amount of limit diffusible hydrogen using the above samples. Test No. in Table 1 1-32 are examples of the present invention. 33 to 40 are comparative examples. As can be seen from the table, all of the examples of the present invention realize a high-strength steel of 1200 to 1600 MPa with excellent delayed fracture resistance with a hydrogen trap amount of 2 ppm or more and a limit diffusible hydrogen amount of 5 ppm or more. Has been.
これに対して比較例を見ると、まず、試験Ｎｏ．３３は炭素量が少ないために、１２００ＭＰａの引張強さに到達していない例、逆に試験Ｎｏ．３４は炭素量が多いために、強度が１６００ＭＰａを越え、延性が不足する例である。 On the other hand, looking at the comparative example, first, test No. No. 33 is an example in which the tensile strength of 1200 MPa has not been reached because of the small amount of carbon. No. 34 is an example in which the strength exceeds 1600 MPa and the ductility is insufficient due to the large amount of carbon.
試験Ｎｏ．３４〜３６は合金−Ｃの質量％比である（０．１８Ｖ＋０．０６Ｍｏ＋０．２５Ｔｉ＋０．１３Ｎｂ）／Ｃが０．５未満であるために、十分な水素トラップ量、限界拡散性水素量が得られなかった例である。 Test No. Since 34 to 36 is a mass% ratio of alloy-C (0.18V + 0.06Mo + 0.25Ti + 0.13Nb) / C is less than 0.5, a sufficient amount of hydrogen trap and limit diffusible hydrogen can be obtained. This is an example that did not exist.
試験Ｎｏ．３７、３８は溶体化温度が低かったため、粗大な未溶解炭化物が残留し、微細合金炭化物の析出がなく、水素トラップ量、限界拡散性水素量が見られず、強度も１２００ＭＰａの引張強さに到達していない例である。 Test No. 37 and 38 had a low solution temperature, coarse undissolved carbide remained, no precipitation of fine alloy carbide, no hydrogen trap amount, no limit diffusible hydrogen amount, and a tensile strength of 1200 MPa. This is an example that has not been reached.
試験Ｎｏ．３９、４０は焼戻し温度が５５０℃未満であるか、または７００℃を越えるため、十分な水素トラップ量、限界拡散性水素量または引張強さが得られなかった例である。 Test No. Nos. 39 and 40 are examples in which a sufficient amount of hydrogen trap, limit diffusible hydrogen or tensile strength cannot be obtained because the tempering temperature is less than 550 ° C. or exceeds 700 ° C.
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| US10060015B2 (en) | 2013-09-25 | 2018-08-28 | Kobe Steel, Ltd. | Steel for high-strength bolts which has excellent delayed fracture resistance and bolt formability, and bolt |
| EP3050987A4 (en) * | 2013-09-25 | 2017-05-31 | Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) | Steel for high-strength bolts which has excellent delayed fracture resistance and bolt moldability, and bolt |
| WO2016031528A1 (en) * | 2014-08-29 | 2016-03-03 | 日産自動車株式会社 | Steel for high-strength bolt, and high-strength bolt |
| JP2016050329A (en) * | 2014-08-29 | 2016-04-11 | 日産自動車株式会社 | Steel for high strength bolt and high strength bolt |
| WO2018235962A1 (en) | 2017-06-23 | 2018-12-27 | 新日鐵住金株式会社 | High-strength steel member |
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| CN109161794B (en) * | 2018-08-31 | 2019-11-22 | 钢铁研究总院 | A kind of anti-delayed fracture high strength steel and preparation method thereof |
| KR20210148334A (en) | 2019-05-14 | 2021-12-07 | 닛폰세이테츠 가부시키가이샤 | Bolts, and steel for bolts |
| JP7464832B2 (en) | 2019-05-14 | 2024-04-10 | 日本製鉄株式会社 | Bolts and bolt steel |
| JPWO2021070384A1 (en) * | 2019-10-11 | 2021-04-15 | ||
| WO2021070384A1 (en) * | 2019-10-11 | 2021-04-15 | 日本製鉄株式会社 | High-strength steel member |
| JP2022553264A (en) * | 2019-11-11 | 2022-12-22 | ロベルト・ボッシュ・ゲゼルシャフト・ミト・ベシュレンクテル・ハフツング | A slow transformation steel alloy, a method of producing a slow transformation steel alloy, and a hydrogen reservoir having components made of the slow transformation steel alloy |
| CN115404399A (en) * | 2021-05-28 | 2022-11-29 | 宝山钢铁股份有限公司 | Homogeneous high-strength durable bolt steel and preparation method thereof |
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