JP4448456B2 - Case-hardened steel with excellent coarse grain prevention and fatigue characteristics during carburizing and its manufacturing method - Google Patents
Case-hardened steel with excellent coarse grain prevention and fatigue characteristics during carburizing and its manufacturing method Download PDFInfo
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本発明は、浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼とその製造方法に関するものである。 The present invention relates to a case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing, and a method for producing the same.
歯車、軸受部品、転動部品、シャフト、等速ジョイント部品は、通常、例えばＪＩＳ Ｇ ４０５２、ＪＩＳ Ｇ ４１０４、ＪＩＳ Ｇ ４１０５、ＪＩＳ Ｇ ４１０６などに規定されている中炭素の機械構造用合金鋼を使用し、冷間鍛造（転造も含む）又は熱間鍛造−切削により所定の形状に加工された後、浸炭焼入れを行う工程で製造されている。冷間鍛造は、製品の表面肌、寸法精度が良く、熱間鍛造に比べて製造コストが低く、歩留まりも良好であるため、従来は熱間鍛造で製造されていた部品を、冷間鍛造へ切り替える傾向が強くなっており、冷鍛−浸炭工程で製造される浸炭部品の対象は近年顕著に増加している。浸炭部品の大きな課題として、熱処理歪みの低減が挙げられる。これは、シャフトについては熱処理歪みで曲がればシャフトとしての機能が損なわれるためであり、また歯車や等速ジョイント部品では熱処理歪みが大きければ、騒音や振動の原因となるためである。ここで、熱処理歪みの最大の原因は、浸炭時に発生する粗大粒である。この粗大粒を抑制するために、従来は、冷間鍛造後、浸炭焼入れの前に、焼鈍が行われていた。これに対して、コスト削減の視点から、近年焼鈍省略の指向が強い。そのためには、焼鈍を省略しても粗大粒を生じない鋼材が強く求められている。
一方、歯車、軸受部品、転動部品のなかで高面圧が負荷される軸受部品、転動部品においては、高深度浸炭が行われている。高深度浸炭は、通常、十数時間から数十時間の長時間を要するために、省エネルギーの視点から、浸炭時間の短縮が重要な課題である。浸炭時間短縮のためには、浸炭温度の高温化が有効である。通常の浸炭温度は９３０℃程度であるが、これに対して、９９０〜１０９０℃の温度域でいわゆる高温浸炭を行うと、粗大粒が発生し、必要な疲労特性、転動疲労特性等が得られないという問題が発生している。そのため、高温浸炭でも粗大粒が発生しない、つまり高温浸炭に適した肌焼き鋼が求められている。このような高面圧が負荷される歯車、軸受部品、転動部品は大型部品が多く、通常「棒鋼−熱間鍛造−必要により焼準等の熱処理−切削−浸炭焼入れ−必要により研磨」の工程で製造される。浸炭時の粗大粒の発生を抑制するためには、熱間鍛造後の状態で、つまり熱間鍛造部材の状態で、粗大粒を抑制するために適正な材質を造り込んでおくことが必要があるが、そのためには、棒鋼線材の素材の状態で粗大粒を抑制するために適正な材質を造り込んでおくことが必要がある。
Gears, bearing parts, rolling parts, shafts, constant velocity joint parts are usually made of medium carbon alloy steel for machine structure as defined in JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc. It is manufactured in a process of carburizing and quenching after being processed into a predetermined shape by cold forging (including rolling) or hot forging-cutting. Cold forging has good surface texture and dimensional accuracy of the product, has a lower manufacturing cost than hot forging, and has a good yield, so parts that were conventionally manufactured by hot forging are now cold forged. The tendency to switch is increasing, and the number of carburized parts manufactured in the cold forging-carburizing process has increased significantly in recent years. A major problem with carburized parts is the reduction of heat treatment distortion. This is because if the shaft is bent by heat treatment distortion, the function as the shaft is impaired, and if the heat treatment distortion is large in gears or constant velocity joint parts, noise and vibration are caused. Here, the largest cause of heat treatment distortion is coarse particles generated during carburization. In order to suppress the coarse grains, conventionally, annealing is performed after cold forging and before carburizing and quenching. On the other hand, from the viewpoint of cost reduction, there is a strong tendency to omit annealing in recent years. For this purpose, there is a strong demand for steel materials that do not produce coarse grains even if annealing is omitted.
On the other hand, high-depth carburization is performed on bearing parts and rolling parts to which high surface pressure is applied among gears, bearing parts, and rolling parts. Since deep carburization usually requires a long time of 10 to several tens of hours, shortening the carburizing time is an important issue from the viewpoint of energy saving. Increasing the carburizing temperature is effective for shortening the carburizing time. The normal carburizing temperature is about 930 ° C. On the other hand, when so-called high-temperature carburizing is performed in the temperature range of 990 to 1090 ° C, coarse grains are generated, and necessary fatigue characteristics, rolling fatigue characteristics, etc. are obtained. The problem of not being able to occur. Therefore, there is a demand for case-hardened steel that does not generate coarse grains even at high temperature carburization, that is, suitable for high temperature carburization. Gears, bearing parts, and rolling parts that are loaded with such high surface pressures are often large parts, and are usually "bar steel-hot forging-heat treatment such as normalization if necessary-cutting-carburizing and quenching-polishing if necessary". Manufactured in a process. In order to suppress the generation of coarse grains during carburizing, it is necessary to build in an appropriate material in order to suppress coarse grains in the state after hot forging, that is, in the state of a hot forged member. However, for that purpose, it is necessary to build in an appropriate material in order to suppress coarse grains in the state of the material of the steel bar wire.
従来技術には、特定量のＡｌ、Ｎを含有し、熱間圧延方向に平行な断面の組織のフェライトバンドの状態を適正化した粗大粒防止特性に優れた肌焼鋼が開示されている（例えば、特許文献１参照）。しかしながら、この肌焼鋼は、球状化焼鈍−冷間鍛造工程を経て製造される部品については粗大粒抑制の能力は不安定であり、また高温浸炭においても粗大粒の発生を抑制できない場合があるのが現実である。 In the prior art, a case-hardened steel containing a specific amount of Al and N and having an excellent coarse grain prevention characteristic in which the state of the ferrite band of the cross-sectional structure parallel to the hot rolling direction is optimized is disclosed ( For example, see Patent Document 1). However, this case-hardened steel has an unstable ability to suppress coarse grains for parts manufactured through the spheroidizing annealing-cold forging process, and it may not be possible to suppress the occurrence of coarse grains even in high-temperature carburizing. Is the reality.
また、Ｔｉ：０．１０〜０．３０％、Ｎ：０．０１％以下、鋼片熱間圧延加熱：１２５０〜１４００℃、製品圧延加熱：Ａｃ３〜１０５０℃を特徴とする疲労特性に優れた肌焼鋼および肌焼鋼の製造方法が開示されている（例えば、特許文献２及び３参照）。特許文献２の実施例で開示されているＮ量は０．００６４〜０．００９６％、特許文献３の実施例で開示されているＮ量は０．００５５〜０．００８４％の範囲である。本鋼もやはり、粗大粒抑制の能力は不安定であり、粗大粒の発生を抑制できる場合もあれば、できない場合もある。また、該鋼は上記の実施例から明らかな通り、０．００５５％以上のＮを含有するとともに、製品圧延加熱：Ａｃ３〜１０５０℃のために、ＴｉＮが粗大化する。そのため、この粗大なＴｉＮが転動疲労やピッチング疲労の起点となるために、十分な疲労特性が得られないという欠点を有している。 Also, excellent fatigue characteristics characterized by Ti: 0.10 to 0.30%, N: 0.01% or less, steel slab hot rolling heating: 1250 to 1400 ° C., product rolling heating: Ac 3 to 1050 ° C. Case-hardened steel and a method for producing case-hardened steel are disclosed (for example, see Patent Documents 2 and 3). The amount of N disclosed in the example of Patent Document 2 is 0.0064 to 0.0096%, and the amount of N disclosed in the example of Patent Document 3 is in the range of 0.0055 to 0.0084%. Again, this steel has an unstable ability to suppress coarse grains and may or may not be able to suppress the generation of coarse grains. Further, as apparent from the above examples, the steel contains 0.0055% or more of N, and TiN coarsens due to product rolling heating: Ac3 to 1050 ° C. For this reason, this coarse TiN serves as a starting point for rolling fatigue and pitching fatigue, and thus has a drawback that sufficient fatigue characteristics cannot be obtained.
さらに、Ｔｉ：０．１超〜０．２％、Ｎ：０．０１５％以下、旧オーステナイト結晶粒度が１１番以上まで微細化したマルテンサイト組織からなる高強度肌焼鋼、およびＮ：０．０２０％以下で「Ｔｉ：０．０５〜０．２％、Ｖ：０．０２〜０．１０％、Ｎｂ：０．０２〜０．１％」のうち１種ないし２種以上を含有し、旧オーステナイト結晶粒度が１１番以上まで微細化したマルテンサイト組織からなる高強度肌焼鋼が開示されている（例えば、特許文献４参照）。本公報の実施例で開示されているＮ量は０．００５２〜０．００９３％である。また、本公報において「旧オーステナイト結晶粒度が１１番以上」としているのは、本公報明細書０００６、０００７および実施例の記載から、鋼素材の状態ではなく、鋼部品の最終的な熱処理を終えた状態において観察される旧オーステナイト結晶粒度である。そして、本公報において「旧オーステナイト結晶粒度が１１番以上」を達成するための方法は、本公報明細書０００８、０００９、００１０および実施例００１１の記載から、浸炭焼入れした後、再焼入れ（再加熱してオーステナイト化した後再び焼入れする方法）することを前提としている。再加熱焼入れを行なうと、熱処理変形が大きくなり強度特性が劣化するという欠点を有する。また、「旧オーステナイト結晶粒度が１１番以上」まで微細化すると、焼入れ性の低下が顕著になって十分な硬さが得られず、この点からも、かえって強度特性の低下を引起す。さらに、該鋼は上記の実施例から明らかな通り、０．００５２％以上のＮを含有するために粗大なＴｉＮが生成する。そのため、この粗大なＴｉＮが転動疲労やピッチング疲労の起点となるために、十分な疲労特性が得られないという欠点を有している。
上記のような開示された方法では、浸炭焼入れ工程において粗大粒の発生を安定的に抑制することができず、歪みや曲がりの発生を安定的に防止することはできない。また、疲労特性、特に転動疲労特性が要求される軸受部品、転動部品についても、高温浸炭により高深度浸炭を行って、十分な疲労特性を実現するのは困難である。本発明はこのような問題を解決して、熱処理歪みの小さい浸炭時の粗大粒防止特性と疲労特性に優れた肌焼き鋼とその製造方法を提供するものである。 In the disclosed method as described above, the generation of coarse particles cannot be stably suppressed in the carburizing and quenching process, and the generation of distortion and bending cannot be stably prevented. In addition, it is difficult to achieve sufficient fatigue characteristics by performing high-depth carburization by high-temperature carburization for bearing parts and rolling parts that require fatigue characteristics, particularly rolling fatigue characteristics. The present invention solves such problems, and provides a case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing with a small heat treatment strain and a method for producing the same.
本発明者らは、上記目的を達成するために結晶粒の粗大化の支配因子について鋭意調査し、次の点を明らかにした。 In order to achieve the above-mentioned object, the present inventors diligently investigated the governing factor of the coarsening of crystal grains and clarified the following points.
（１） 浸炭時に結晶粒の粗大化を防止するには、ピン止め粒子としてＡｌＮ、ＮｂＮを活用するよりも、ＴｉＣ、ＴｉＣＳを主体とするＴｉ系析出物を浸炭時に微細析出させることが有効である。またはさらに上記と併用して、ＮｂＣを主体とするＮｂの炭窒化物を浸炭時に微細析出させることにより、粗大粒防止特性は一層向上する。
（２） 上記のＴｉ系析出物またはさらにＮｂの炭窒化物を浸炭時に微細析出させる方法として、新規に以下の方法を発見した。
（ａ）上記のＴｉ系析出物またはさらにＮｂの炭窒化物を浸炭時にピン止め粒子として活用するためには、浸炭焼入れ時にこれらの析出物を多量微細分散する必要がある。そのためには、棒鋼または線材を熱間圧延する場合の圧延加熱時に上記のＴｉ系析出物またはさらにＮｂの炭窒化物を一旦溶体化する必要がある。Ｎ量が高くて、圧延加熱時にＴｉＮが多量に残存すると、Ｔｉ系析出物はＴｉＮ主体の複合析出物となり、溶体化が困難となる。また、熱間圧延後の冷却過程で、粗大なＴｉＮ上にＴｉＣ、ＴｉＣＳまたはさらにＮｂＣが析出し、これらの析出物の微細分散が妨げられる。そのため、Ｎ量を出来るだけ低減することが必要である。
（ｂ）また、圧延加熱時に粗大なＡｌＮが存在すると、上記のＴｉＮと同じ悪影響を及ぼす。そのため、ＡｌＮも圧延加熱時に溶体化しておく必要がある。ここで、ＡｌＮを圧延加熱時に溶体化しておけば、棒鋼、線材の熱間圧延−冷却過程でＡｌＮの析出はほとんど起こらない。そのため、熱間圧延後のＡｌＮの析出量を規制することにより、圧延加熱時のＡｌＮの溶体化状況の確認が可能である。
（ｃ）なお、ＡｌＮが圧延加熱時に溶体化できる条件で加熱を行えば、Ｔｉ系析出物またはさらにＮｂＣの析出物を一旦溶体化することが可能である。そのため、熱間圧延後のＡｌＮの析出量を規制することにより、Ｔｉ系析出物またはさらにＮｂＣの析出物を圧延加熱時に一旦溶体化できたことの確認が可能である。
（ｄ）さらに、Ｔｉ系析出物またはさらにＮｂＣの析出物のピン止め効果を安定して発揮させるには、熱間圧延後のマトリックス中にこれらの析出物を微細析出させておくことが必要である。そのためには、熱間圧延時の冷却過程でオーステナイトからの拡散変態時に相界面析出させる必要がある。もし熱間圧延ままの組織にベイナイトが生成すると、上記の析出物の相界面析出が困難になるために、ベイナイトを実質的に含まない組織とすることが必須である。
(1) In order to prevent coarsening of crystal grains during carburizing, it is effective to finely precipitate Ti-based precipitates mainly composed of TiC and TiCS during carburizing rather than using AlN and NbN as pinning particles. is there. Or, in combination with the above, the coarse grain prevention characteristics are further improved by finely precipitating Nb carbonitrides mainly composed of NbC during carburizing.
(2) The following method was newly discovered as a method for finely precipitating the Ti-based precipitate or the Nb carbonitride during carburization.
(A) In order to utilize the above Ti-based precipitates or Nb carbonitrides as pinning particles during carburizing, it is necessary to finely disperse these precipitates in a large amount during carburizing and quenching. For this purpose, the Ti-based precipitate or the Nb carbonitride needs to be once solutionized at the time of rolling and heating when hot rolling a steel bar or wire. If the amount of N is high and a large amount of TiN remains during rolling and heating, the Ti-based precipitate becomes a composite precipitate mainly composed of TiN, which makes it difficult to form a solution. Further, in the cooling process after hot rolling, TiC, TiCS or further NbC is precipitated on coarse TiN, and the fine dispersion of these precipitates is hindered. Therefore, it is necessary to reduce the N amount as much as possible.
(B) In addition, if coarse AlN is present during rolling and heating, the same adverse effect as TiN described above. Therefore, it is necessary to make AlN into solution during rolling and heating. Here, if AlN is formed into a solution at the time of rolling and heating, precipitation of AlN hardly occurs in the hot rolling-cooling process of the steel bar and the wire. Therefore, by regulating the precipitation amount of AlN after hot rolling, it is possible to confirm the solution state of AlN during rolling and heating.
(C) If heating is performed under the condition that AlN can be dissolved during rolling and heating, Ti-based precipitates or further NbC precipitates can be once solutionized. Therefore, by regulating the amount of AlN deposited after hot rolling, it is possible to confirm that Ti-based precipitates or further NbC precipitates were once formed into a solution during rolling and heating.
(D) Furthermore, in order to stably exhibit the pinning effect of Ti-based precipitates or further NbC precipitates, it is necessary to finely precipitate these precipitates in the matrix after hot rolling. is there. For this purpose, it is necessary to cause phase interface precipitation during the diffusion transformation from austenite during the cooling process during hot rolling. If bainite is generated in a structure as hot-rolled, precipitation of the above-described precipitate at the phase interface becomes difficult. Therefore, it is essential to make the structure substantially free of bainite.
（３） 熱間圧延後の鋼材の状態で、ＡｌＮの析出量を極力制限するためには、つまり熱間圧延加熱時にＴｉ系析出物またはさらにＮｂＣの析出物を溶体化するためには、圧延加熱温度を高温にする必要がある。 (3) In order to limit the precipitation amount of AlN as much as possible in the state of the steel material after hot rolling, that is, in order to solutionize Ti-based precipitates or further NbC precipitates during hot rolling heating, rolling It is necessary to increase the heating temperature.
（４） 熱間圧延後の鋼材に、Ｔｉ系析出物またはさらにＮｂＣの析出物をあらかじめ微細析出させるためには、圧延加熱温度及び圧延後の冷却条件を最適化すれば良い。すなわち圧延加熱温度を高温にすることによって、Ｔｉ系析出物またはさらにＮｂＣの析出物を一旦マトリックス中に固溶させ、熱間圧延後にＴｉ系析出物またはさらにＮｂＣの析出物の析出温度域を徐冷することによって、これらの炭窒化物を多量、微細分散させることができる。
（５） さらに、熱間圧延後の鋼材のフェライト粒が過度に微細であると、浸炭加熱時に粗大粒が発生しやすくなるため、圧延仕上げ温度の適正化もポイントである。
（６） Ｔｉ添加鋼ではＴｉ析出物が疲労破壊の起点となるため、疲労特性、特に転動疲労特性が劣化しやすくなるが、低Ｎ化、熱間圧延温度の高温化等によりＴｉ析出物最大サイズを小さくすることにより疲労特性の改善が可能となり、粗大粒防止特性と疲労特性の両立が可能である。
(4) In order to finely precipitate Ti-based precipitates or further NbC precipitates in the steel after hot rolling, the rolling heating temperature and the cooling conditions after rolling may be optimized. That is, by increasing the rolling heating temperature, Ti-based precipitates or further NbC precipitates are once dissolved in the matrix, and after hot rolling, the precipitation temperature range of Ti-based precipitates or further NbC precipitates is gradually increased. By cooling, a large amount of these carbonitrides can be finely dispersed.
(5) Furthermore, if the ferrite grains of the steel material after hot rolling are excessively fine, coarse grains are likely to be generated during carburizing heating, so that optimization of the rolling finishing temperature is also a point.
(6) In Ti-added steels, Ti precipitates are the starting point for fatigue failure, so fatigue characteristics, particularly rolling fatigue characteristics, are likely to deteriorate. However, Ti precipitates are reduced due to lower N and higher hot rolling temperature. By reducing the maximum size, fatigue characteristics can be improved, and both coarse grain prevention characteristics and fatigue characteristics can be achieved.
本発明は以上の新規なる知見にもとづいてなされたものであり、本発明の要旨は以下の通りである。 The present invention has been made based on the above novel findings, and the gist of the present invention is as follows.
（１） 質量％として、
Ｃ：０．１〜０．４％、
Ｓｉ：０．０２〜１．５％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．００５〜０．０５％
Ｔｉ：０．０５〜０．２％、
Ｎ：０．００５１％未満に制限し、
Ｃｒ：０．４〜２．０％、
Ｐ：０．０２５％以下に制限し、
Ｏ：０．００２５％以下に制限し、
残部が鉄および不可避的不純物からなり、
熱間圧延後のＡｌＮの析出量を０．０１％以下に制限したことを特徴とする浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
（２） 鋼が、さらに、質量％で、Ｎｂ：０．０４％未満を含有することを特徴とする上記（１）に記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
（３） 鋼が、さらに、質量％で、Ｍｏ：１．５％以下、Ｎｉ：３．５％以下、V：０．５％以下、及びＢ：０．００５％以下のうちの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）または（２）に記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
（４） 熱間圧延後のベイナイトの組織分率を３０％以下に制限したことを特徴とする上記（１）〜（３）の内のいずれかに記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
（５） 熱間圧延後のフェライト結晶粒度番号がＪＩＳ Ｇ０５５２で規定されている８〜１１番であることを特徴とする上記（１）〜（４）の内のいずれかに記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
（６） 熱間圧延後の鋼のマトリックス中の長手方向断面において、検査基準面積：１００平方ｍｍ、検査回数：１６視野、予測を行なう面積：３００００平方ｍｍの条件で測定された極値統計によるＴｉ系析出物の最大直径が４０μｍ以下であることを特徴とする上記（１）〜（５）の内のいずれかに記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。
(1) As mass%,
C: 0.1-0.4%
Si: 0.02 to 1.5%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.005 to 0.05%
Ti: 0.05-0.2 %
N : limited to less than 0.0051% ,
Cr : 0.4 to 2.0 %,
P : limited to 0.025% or less,
O: limited to 0.0025% or less,
The balance consists of iron and inevitable impurities,
A case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing, characterized by limiting the precipitation amount of AlN after hot rolling to 0.01% or less .
(2) Case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing according to (1) above, wherein the steel further contains, by mass%, Nb: less than 0.04% .
(3) The steel is further in mass%, Mo: 1.5% or less, Ni: 3.5% or less, V: 0.5% or less, and B: 0.005% or less, or Case hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburization according to the above (1) or (2), characterized by containing two or more.
(4) The coarse grain prevention property and fatigue during carburizing according to any one of the above (1) to (3), wherein the structure fraction of bainite after hot rolling is limited to 30% or less Case-hardened steel with excellent properties.
(5) The ferrite grain size number after hot rolling is No. 8 to No. 11 defined in JIS G0552, and the carburizing time according to any one of (1) to (4) above Case-hardened steel with excellent coarse grain prevention and fatigue characteristics.
(6) According to extreme value statistics measured under conditions of inspection standard area: 100 square mm, number of inspections: 16 fields of view, prediction area: 30000 square mm, in longitudinal section in steel matrix after hot rolling The maximum diameter of the Ti-based precipitate is 40 μm or less, and the case hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing according to any one of the above (1) to (5).
（７） 質量％として、 (7) As mass%,
Ｃ：０．１〜０．４％、C: 0.1-0.4%
Ｓｉ：０．０２〜１．５％、Si: 0.02 to 1.5%,
Ｍｎ：０．３〜１．８％、Mn: 0.3 to 1.8%
Ｓ：０．００１〜０．１５％、S: 0.001 to 0.15%,
Ａｌ：０．００５〜０．０５％Al: 0.005 to 0.05%
Ｔｉ：０．０５〜０．２％、Ti: 0.05 to 0.2%,
Ｎ：０．００５１％未満に制限し、N: limited to less than 0.0051%,
Ｃｒ：０．４〜２．０％、Cr: 0.4 to 2.0%,
Ｐ：０．０２５％以下に制限し、P: limited to 0.025% or less,
Ｏ：０．００２５％以下に制限し、O: limited to 0.0025% or less,
残部が鉄および不可避的不純物からなる鋼を、Steel with the balance of iron and inevitable impurities
１１５０℃以上の温度で保熱時間１０分以上加熱して線材または棒鋼に熱間圧延し、熱間圧延後のＡｌＮの析出量を０．０１％以下に制限した鋼となるようにすることを特徴とする浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼の製造方法。Heating at a temperature of 1150 ° C. or higher for 10 minutes or longer and hot rolling to a wire or steel bar, so that the precipitation amount of AlN after hot rolling is limited to 0.01% or less. A method for producing case-hardened steel with excellent coarse grain prevention characteristics and fatigue characteristics during carburizing.
（８） 鋼が、さらに、質量％でＮｂ：０．０４％未満を含有することを特徴とする上記（７）に記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼の製造方法。 (8) The steel further contains Nb: less than 0.04% by mass%, and the case-hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing according to (7) above Production method.
（９） 鋼が、さらに、質量％で、Ｍｏ：１．５％以下、Ｎｉ：３．５％以下、V：０．５％以下、及びＢ：０．００５％以下のうちの１種または２種以上を含有することを特徴とする上記（７）または（８）に記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼の製造方法。 (9) The steel is further in mass%, Mo: 1.5% or less, Ni: 3.5% or less, V: 0.5% or less, and B: 0.005% or less, or The method for producing a case hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics at the time of carburizing according to the above (7) or (8), comprising two or more kinds.
（１０） 熱間圧延後に８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷し、熱間圧延後のベイナイトの組織分率が３０％以下の鋼となるようにすることを特徴とする浸炭時の粗大粒防止特性と疲労特性に優れた上記（７）〜（９）の内のいずれかに記載の肌焼鋼の製造方法。 (10) After hot rolling, the temperature range of 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less so that the steel has a bainite structure fraction of 30% or less after hot rolling. The manufacturing method of the case hardening steel in any one of said (7)-(9) excellent in the coarse grain prevention characteristic and the fatigue characteristic at the time of carburizing characterized by these.
（１１） 熱間圧延の仕上げ温度を８４０〜１０００℃とし、フェライト結晶粒度番号がＪＩＳ Ｇ０５５２で規定されている８〜１１番である鋼となるようにすることを特徴とする上記（７）〜（１０）の内のいずれかに記載の浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼の製造方法。 (11) The above-mentioned (7) to (7), characterized in that the finishing temperature of hot rolling is 840 to 1000 ° C., and the steel has a ferrite grain size number of 8 to 11 as defined in JIS G0552. (10) The manufacturing method of the case hardening steel excellent in the coarse grain prevention characteristic and fatigue characteristic at the time of carburizing in any one of.
本発明の浸炭時の粗大粒防止特性に優れた肌焼鋼及びその製造方法を用いれば、冷鍛工程で部品を製造しても、浸炭時の結晶粒の粗大化が抑制されるために、疲労強度特性も優れるとともに、焼入れ歪みによる寸法精度の劣化が従来よりも極めて少ない。そのため、これまで、粗大粒の問題から冷鍛化が困難であった部品の冷鍛化が可能になり、さらに冷鍛後の焼鈍を省略することも可能になる。また、熱間鍛造工程で製造される部品に本鋼材を適用しても高温浸炭においても粗大粒の発生を防止し、転動疲労特性等の十分な強度特性を得ることができる。以上のように、本発明による産業上の効果は極めて顕著なるものがある。 If the case-hardened steel excellent in the coarse grain prevention property during carburizing and the manufacturing method thereof according to the present invention are used, even if parts are manufactured in the cold forging process, the coarsening of crystal grains during carburizing is suppressed. The fatigue strength characteristics are excellent, and the deterioration of dimensional accuracy due to quenching strain is extremely less than in the past. Therefore, it becomes possible to cold forge parts that have been difficult to cold forge due to the problem of coarse grains, and it is also possible to omit annealing after cold forging. Moreover, even if this steel material is applied to a part manufactured in the hot forging process, generation of coarse grains can be prevented even in high-temperature carburization, and sufficient strength characteristics such as rolling fatigue characteristics can be obtained. As described above, the industrial effects of the present invention are extremely remarkable.
以下、本発明について詳細に説明する。 Hereinafter, the present invention will be described in detail.
まず、本発明の肌焼鋼の成分を限定した理由について説明する。 First, the reason which limited the component of the case hardening steel of this invention is demonstrated.
Ｃは鋼に必要な強度を与えるのに有効な元素であるが、０．１％未満では必要な引張強さを確保することができず、０．４％を越えると硬くなって冷間加工性が劣化するとともに、浸炭後の芯部靭性が劣化するので、０．１〜０．４％の範囲内にする必要がある。 C is an element effective for imparting the necessary strength to steel, but if it is less than 0.1%, the required tensile strength cannot be secured, and if it exceeds 0.4%, it becomes hard and cold work is performed. The core portion toughness after carburizing deteriorates as well as the properties deteriorate, so it is necessary to set the content within the range of 0.1 to 0.4%.
Ｓｉは鋼の脱酸に有効な元素であるとともに、鋼に必要な強度、焼入れ性を与え、焼戻し軟化抵抗を向上するのに有効な元素であるが、０．０２％未満ではその効果は不十分である。一方、１．５％を越えると、硬さの上昇を招き冷間鍛造性が劣化する。以上の理由から、その含有量を０．０２〜１．５％の範囲内にする必要がある。冷間加工を受ける鋼材の好適範囲は０．０２〜０．３％である。特に冷鍛性を重視する場合は、０．０２〜０．１５％の範囲にするのが望ましい。一方、Ｓｉは粒界強度の増加に有効な元素であり、さらに軸受部品、転動部品においては、転動疲労過程での組織変化、材質劣化の抑制による高寿命化に有効な元素である。そのため、高強度化を指向する場合には、０．２〜１．５％の範囲が好適である。特に転動疲労強度の高いレベルを得るためには、０．４〜１．５％の範囲にするのが望ましい。 なお、Ｓｉ添加による軸受部品、転動部品の転動疲労過程での組織変化、材質劣化の抑制の効果は、浸炭後の組織中の残留オーステナイト量（通称、残留γ量）が３０〜４０％の時に特に大きい。残留γ量をこの範囲で制御するには、いわゆる浸炭浸窒処理を行うことが有効である。浸炭浸窒処理は、浸炭後の拡散処理の過程で浸窒を行う処理である。表面の窒素濃度が０．２〜０．６％の範囲になるような条件が適切である。なお、この場合の浸炭時の炭素ポテンシャルは０．９〜１．３％の範囲とするのが望ましい。 Si is an element effective for deoxidation of steel, and is an element effective for imparting necessary strength and hardenability to steel and improving temper softening resistance. However, if it is less than 0.02%, the effect is ineffective. It is enough. On the other hand, if it exceeds 1.5%, the hardness is increased and the cold forgeability is deteriorated. For the above reasons, the content needs to be in the range of 0.02 to 1.5%. The suitable range of steel materials that undergo cold working is 0.02 to 0.3%. In particular, when emphasizing cold forgeability, it is desirable to make the range 0.02 to 0.15%. On the other hand, Si is an element effective for increasing the grain boundary strength. Further, in bearing parts and rolling parts, it is an element effective for extending the life by suppressing structural changes and material deterioration during the rolling fatigue process. Therefore, when aiming at high intensity | strength, the range of 0.2 to 1.5% is suitable. In particular, in order to obtain a high level of rolling fatigue strength, it is desirable to be in the range of 0.4 to 1.5%. In addition, the effect of suppressing the structural change and material deterioration in rolling fatigue process of bearing parts and rolling parts due to the addition of Si is 30 to 40% of the retained austenite amount (common name, residual γ amount) in the structure after carburizing. Especially at times. In order to control the residual γ amount within this range, it is effective to perform a so-called carburizing and nitriding treatment. The carburizing and nitriding process is a process of performing nitriding in the process of diffusion after carburizing. Conditions under which the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. In this case, it is desirable that the carbon potential at the time of carburizing is in the range of 0.9 to 1.3%.
Ｍｎは鋼の脱酸に有効な元素であるとともに、鋼に必要な強度、焼入れ性を与えるのに有効な元素であるが、０．３％未満では効果は不十分であり、１．８％を越えるとその効果は飽和するのみならず、硬さの上昇を招き冷間鍛造性が劣化するので、０．３％〜１．８％の範囲内にする必要がある。好適範囲は０．５〜１．２％である。なお、冷鍛性を重視する場合は、０．５〜０．７５％の範囲にするのが望ましい。 Mn is an element effective for deoxidation of steel and is an element effective for imparting the necessary strength and hardenability to the steel, but if less than 0.3%, the effect is insufficient, 1.8% If it exceeds, the effect is not only saturated, but also the hardness is increased and the cold forgeability is deteriorated, so it is necessary to be within the range of 0.3% to 1.8%. The preferred range is 0.5-1.2%. In addition, when importance is attached to cold forgeability, it is desirable to set it as 0.5 to 0.75% of range.
Ｓは鋼中でＭｎＳを形成し、これによる被削性の向上を目的として添加するが、０．００１％未満ではその効果は不十分である。一方、０．１５％を超えるとその効果は飽和し、むしろ粒界偏析を起こし粒界脆化を招く。以上の理由から、Ｓの含有量を０．００１〜０．１５％の範囲内にする必要がある。なお、軸受部品、転動部品においては、ＭｎＳが転動疲労寿命を劣化させるために、Ｓを極力低減する必要があり、０．００１〜０．０１％の範囲にするのが望ましい。 S forms MnS in the steel and is added for the purpose of improving the machinability. However, if it is less than 0.001%, its effect is insufficient. On the other hand, if it exceeds 0.15%, the effect is saturated, and rather, grain boundary segregation occurs, leading to grain boundary embrittlement. For these reasons, the S content needs to be in the range of 0.001 to 0.15%. In addition, in bearing parts and rolling parts, since MnS deteriorates the rolling fatigue life, it is necessary to reduce S as much as possible, and it is desirable to make it in the range of 0.001 to 0.01%.
Ａｌは脱酸剤として添加する。０．００５％未満ではその効果は不十分である。一方、０．０５％を越えると、ＡｌＮが圧延加熱時に溶体化しないで残存し、ＴｉやＮｂの析出物の析出サイトとなり、これらの析出物の微細分散を阻害し、浸炭時の結晶粒の粗大化を助長する。以上の理由から、その含有量を０．００５〜０．０５％の範囲内にする必要がある。好適範囲は０．０２５〜０．０４％である。
Ｔｉは鋼中で微細なＴｉＣ、ＴｉＣＳを生成させ、これにより浸炭時のγ粒の微細化を図るために添加する。しかしながら、０．０５％未満ではその効果は不十分である。一方、Ｔｉを０．２％を超えて添加すると、ＴｉＣによる析出硬化が顕著になり、冷間加工性が顕著に劣化するとともに、ＴｉＮ主体の析出物が顕著となり転動疲労特性が劣化する。以上の理由から、その含有量を０．０５〜０．２％の範囲内にする必要がある。好適範囲は、０．０５〜０．１％未満である。なお、本発明の鋼および熱間鍛造部材は、浸炭加熱時に侵入してくる炭素および窒素と固溶Ｔｉが反応して、浸炭層に微細なＴｉ（ＣＮ）が多量に析出する。そのために、軸受部品、転動部品においては、これらのＴｉ（ＣＮ）が転動疲労寿命の向上に寄与する。したがって、軸受部品、転動部品において、特に高いレベルの転動疲労寿命を指向する場合には、浸炭時の炭素ポテンシャルを０．９〜１．３％の範囲で高めに設定すること、あるいは、いわゆる浸炭浸窒処理を行うことが有効である。浸炭浸窒処理は、上記のように浸炭後の拡散処理の過程で浸窒を行う処理であるが、表面の窒素濃度が０．２〜０．６％の範囲になるような条件が適切である。
Al is added as a deoxidizer. If it is less than 0.005%, the effect is insufficient. On the other hand, if it exceeds 0.05%, AlN remains without solution during rolling and heating, and becomes a precipitation site for precipitates of Ti and Nb, which inhibits fine dispersion of these precipitates, Promotes coarsening. For the above reasons, the content needs to be in the range of 0.005 to 0.05%. The preferred range is 0.025 to 0.04%.
Ti is added in order to produce fine TiC and TiCS in the steel and thereby refine the γ grains during carburization. However, if it is less than 0.05%, the effect is insufficient. On the other hand, when Ti is added in excess of 0.2%, precipitation hardening due to TiC becomes remarkable, cold workability is remarkably deteriorated, and precipitates mainly composed of TiN become remarkable, and rolling fatigue characteristics deteriorate. For the above reasons, the content needs to be in the range of 0.05 to 0.2%. The preferred range is 0.05 to less than 0.1%. In the steel and hot forged member of the present invention, carbon and nitrogen that enter during carburizing heating react with solute Ti, and a large amount of fine Ti (CN) precipitates in the carburized layer. Therefore, in bearing parts and rolling parts, these Ti (CN) contributes to the improvement of rolling fatigue life. Therefore, in bearing parts and rolling parts, in the case of directing a particularly high level of rolling fatigue life, the carbon potential at the time of carburizing should be set higher in the range of 0.9 to 1.3%, or It is effective to perform so-called carburizing and nitriding treatment. The carburizing and nitriding treatment is a treatment in which nitriding is performed in the course of the diffusion treatment after carburizing as described above, but the conditions that the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. is there.
Ｎは鋼中のＴｉと結びつくと粒制御にほとんど寄与しない粗大なＴｉＮを生成し、これがＴｉＣ、ＴｉＣＳ主体のＴｉ系析出物、ＮｂＣ、ＮｂＣ主体のＮｂ（ＣＮ）の析出サイトとなり、これらのＴｉ系析出物、Ｎｂの炭窒化物の微細析出を阻害し粗大粒の生成を促進する。上記の悪影響はＮ量が０．００５１％以上の場合特に顕著である。以上の理由から、その含有量を０．００５１％未満にする必要がある。 When N is combined with Ti in the steel, coarse TiN that hardly contributes to grain control is generated, and this becomes a precipitation site of Ti-based precipitates mainly composed of TiC and TiCS, NbC and NbC-based Nb (CN), and these Ti Inhibits fine precipitation of system precipitates and Nb carbonitrides and promotes the formation of coarse particles. The above adverse effect is particularly remarkable when the N amount is 0.0051% or more. For the above reasons, the content needs to be less than 0.0051%.
次に、本発明では、Ｃｒを含有する。さらには、Ｍｏ、Ｎｉ、Ｖ、及びＢのうちの１種または２種以上を含有するが好ましい。
Next, in the present invention, Cr is contained. Furthermore, it is preferable to contain one or more of Mo, Ni, V, and B.
Ｃｒは鋼に強度、焼入れ性を与えるのに有効な元素であり、さらに軸受部品、転動部品においては、浸炭後の残留γ量を増大させるとともに、転動疲労過程での組織変化、材質劣化の抑制による高寿命化に有効な元素である。０．４％未満ではその効果は不十分であり、２．０％を超えて添加すると硬さの上昇を招き冷間鍛造性が劣化する。以上の理由から、その含有量を０．４〜２．０％の範囲内にする必要がある。好適範囲は０．７〜１．６％である。なお、Ｃｒ添加による軸受部品、転動部品の転動疲労過程での組織変化、材質劣化の抑制の効果は、浸炭後の組織中の残留γ量が３０〜４０％の時に特に大きい。残留γ量をこの範囲で制御するには、いわゆる浸炭浸窒処理を行い、表面の窒素濃度が０．２〜０．６％の範囲になるようにすることが有効である。 Cr is an element effective for imparting strength and hardenability to steel. Further, in bearing parts and rolling parts, the amount of residual γ after carburizing is increased, and the structure changes and material deterioration occurs during rolling fatigue. It is an element that is effective in extending the life by suppressing the above. If it is less than 0.4%, the effect is insufficient, and if it exceeds 2.0%, the hardness is increased and the cold forgeability deteriorates. For the above reasons, the content needs to be in the range of 0.4 to 2.0%. A preferable range is 0.7 to 1.6%. The effect of suppressing the structural change and material deterioration in the rolling fatigue process of bearing parts and rolling parts due to the addition of Cr is particularly great when the amount of residual γ in the structure after carburizing is 30 to 40%. In order to control the amount of residual γ within this range, it is effective to perform a so-called carburizing and nitriding treatment so that the surface nitrogen concentration is in the range of 0.2 to 0.6%.
Ｍｏも鋼に強度、焼入れ性を与えるのに有効な元素であり、さらに軸受部品、転動部品においては、浸炭後の残留γ量を増大させるとともに、転動疲労過程での組織変化、材質劣化の抑制による高寿命化に有効な元素である。但し、１．５％を超えて添加すると硬さの上昇を招き切削性、冷間鍛造性が劣化する。以上の理由から、その含有量を１．５％以下の範囲内にする必要がある。好ましくは、０．５％以下、さらに好適範囲は０．０２〜０．５％である。Ｍｏ添加による軸受部品、転動部品の転動疲労過程での組織変化、材質劣化の抑制の効果についても、Ｃｒと同様に、いわゆる浸炭浸窒処理を行い、浸炭後の組織中の残留γ量が３０〜４０％の時に特に大きい。 Mo is also an effective element for imparting strength and hardenability to steel. In addition, in bearing parts and rolling parts, the amount of residual γ after carburizing is increased, and the structure changes and material deterioration occurs during rolling fatigue. It is an element that is effective in extending the life by suppressing the above. However, if added over 1.5%, the hardness is increased, and the machinability and cold forgeability deteriorate. For the above reasons, the content needs to be in the range of 1.5% or less. Preferably, it is 0.5% or less, and more preferably 0.02 to 0.5%. As with Cr, the effect of suppressing the structure change and material deterioration during rolling fatigue of bearing parts and rolling parts due to the addition of Mo is the so-called carburizing and nitriding treatment, and the amount of residual γ in the structure after carburizing. Is particularly large at 30-40%.
Ｎｉも鋼に強度、焼入れ性を与えるのに有効な元素であるが、３．５％を超えて添加すると硬さの上昇を招き冷間鍛造性が劣化する。以上の理由から、その含有量を３．５％以下の範囲内にする必要がある。好ましくは０．１〜３．５％、さらに好適範囲は０．４〜２．０％である。なお、Ｎｉ含有量の下限は特に限定するものではないが、０．１％とすることが好ましい。 Ni is also an element effective for imparting strength and hardenability to steel, but if added over 3.5%, the hardness is increased and cold forgeability deteriorates. For the above reasons, the content needs to be in the range of 3.5% or less. Preferably it is 0.1 to 3.5%, and a more preferable range is 0.4 to 2.0%. The lower limit of the Ni content is not particularly limited, but is preferably 0.1%.
Ｖも鋼に強度、焼入れ性を与えるのに有効な元素であるが、０．５％を超えて添加すると硬さの上昇を招き冷間鍛造性が劣化する。以上の理由から、その含有量を０．５％以下の範囲内にする必要がある。好ましくは０．０３〜０．５％、さらに好適範囲は０．０７〜０．２％である。
Ｂも鋼に強度、焼入れ性を与えるのに有効な元素である。さらにＢには、（ａ）棒鋼・線材圧延において、圧延後の冷却過程でボロン鉄炭化物を生成することにより、フェライトの成長速度を増加させ、圧延ままで軟質化を促進する効果、（ｂ）浸炭材の粒界強度を向上させることにより、浸炭部品としての疲労強度・衝撃強度を向上させる効果も有している。しかしながら、０．００５％を超えてＢを添加するとその効果は飽和し、かえって衝撃強度劣化等の悪影響が懸念されるので、その含有量を０．００５％以下の範囲内にする必要がある。好適範囲は０．０００５〜０．００３％である。
V is also an element effective for imparting strength and hardenability to the steel, but if added over 0.5%, the hardness is increased and cold forgeability deteriorates. For the above reasons, the content needs to be in the range of 0.5% or less. Preferably it is 0.03-0.5%, and a more suitable range is 0.07-0.2%.
B is also an element effective for imparting strength and hardenability to steel. Further, in B, (a) the effect of increasing the growth rate of ferrite by generating boron iron carbide in the cooling process after rolling in steel bar / wire rolling, and promoting softening as it is rolled, (b) By improving the grain boundary strength of the carburized material, it also has the effect of improving fatigue strength and impact strength as a carburized component. However, when B is added over 0.005%, the effect is saturated, and adverse effects such as deterioration of impact strength are concerned. Therefore, the content needs to be within the range of 0.005% or less. The preferred range is 0.0005 to 0.003%.
Ｐは冷間鍛造時の変形抵抗を高め、靭性を劣化させる元素であるため、冷間鍛造性が劣化する。また、焼入れ、焼戻し後の部品の結晶粒界を脆化させることによって、疲労強度を劣化させるので、できるだけ低減することが望ましい。従ってその含有量を０．０２５％以下に制限する必要がある。好適範囲は０．０１５％以下である。 Since P is an element that increases deformation resistance during cold forging and deteriorates toughness, cold forgeability deteriorates. Further, since the fatigue strength is deteriorated by embrittlement of the grain boundaries of the parts after quenching and tempering, it is desirable to reduce them as much as possible. Therefore, it is necessary to limit the content to 0.025% or less. The preferred range is 0.015% or less.
本発明のような高Ｔｉ鋼においては、Ｏは鋼中でＴｉ系の酸化物系介在物を形成する。酸化物系介在物が鋼中に多量に存在すると、ＴｉＣの析出サイトとなり、熱間圧延時にＴｉＣが粗大に析出し、浸炭時に結晶粒の粗大化を抑制できなくなる。そのため、Ｏ量はできるだけ低減することが望ましい。以上の理由から、その含有量を０．００２５％以下に制限する必要がある。好適範囲は０．００２０％以下である。なお、軸受部品、転動部品においては、酸化物系介在物が転動疲労破壊の起点となるので、Ｏ含有量が低いほど転動寿命は向上する。そのため、軸受部品、転動部品においては、Ｏ含有量を０．００１２％以下に制限するのが望ましい。 In the high Ti steel as in the present invention, O forms Ti-based oxide inclusions in the steel. When a large amount of oxide inclusions are present in the steel, TiC precipitates, and TiC precipitates coarsely during hot rolling, making it impossible to suppress grain coarsening during carburizing. Therefore, it is desirable to reduce the amount of O as much as possible. For the above reasons, the content needs to be limited to 0.0025% or less. The preferred range is 0.0020% or less. In bearing parts and rolling parts, oxide inclusions are the starting point of rolling fatigue failure, so the rolling life is improved as the O content is lower. Therefore, it is desirable to limit the O content to 0.0012% or less in bearing parts and rolling parts.
次に、本発明では、熱間圧延後のＡｌＮの析出量を０．０１％以下に制限するが、このように限定した理由を以下に述べる。 Next, in the present invention, the precipitation amount of AlN after hot rolling is limited to 0.01% or less. The reason for this limitation will be described below.
圧延加熱時に粗大なＡｌＮが存在すると、Ｔｉ系析出物、Ｎｂの析出物の析出サイトとなり、熱間圧延後にＴｉ系析出物、Ｎｂの析出物が粗大に析出し、浸炭時に結晶粒の粗大化を抑制できなくなる。そのため、圧延加熱時にＡｌＮを溶体化することが必要である。ここで、ＡｌＮは、圧延加熱時に溶体化しておけば、棒鋼、線材の熱間圧延−冷却過程でＡｌＮの析出はほとんど起こらない。そのため、熱間圧延後のＡｌＮの析出量を規制することにより、圧延加熱時にＡｌＮが十分に溶体化できていることの確認が可能である。なお、Ｔｉ系析出物、Ｎｂの析出物をピン止め粒子として活用するためには、圧延加熱時にこれらの析出物も一旦溶体化する必要がある。ＡｌＮが圧延加熱時に溶体化できる条件で加熱を行えば、上記の析出物を一旦溶体化することが可能である。そのため、熱間圧延後のＡｌＮの析出量を規制することにより、圧延加熱時にＴｉ系析出物、Ｎｂの析出物を一旦溶体化できたことの確認が可能である。ＡｌＮの析出量が０．０１％を超えると、粗大粒の発生が懸念される。以上の理由から、熱間圧延後のＡｌＮの析出量を０．０１％以下に制限する。好適範囲は、０．００５％以下である。 When coarse AlN is present during rolling and heating, it becomes a precipitation site for Ti-based precipitates and Nb precipitates, and after hot rolling, Ti-based precipitates and Nb precipitates are coarsely precipitated, and the grains become coarse during carburizing. Can not be suppressed. Therefore, it is necessary to solutionize AlN during rolling and heating. Here, if AlN is in solution at the time of rolling and heating, precipitation of AlN hardly occurs in the hot rolling-cooling process of steel bars and wires. Therefore, it is possible to confirm that AlN can be sufficiently formed into a solution during rolling and heating by regulating the precipitation amount of AlN after hot rolling. In order to utilize Ti-based precipitates and Nb precipitates as pinning particles, these precipitates also need to be dissolved once during rolling and heating. If heating is performed under conditions that allow AlN to form a solution during rolling and heating, the above precipitate can be once formed into a solution. Therefore, by regulating the precipitation amount of AlN after hot rolling, it is possible to confirm that the Ti-based precipitates and Nb precipitates were once in solution during rolling and heating. If the precipitation amount of AlN exceeds 0.01%, the generation of coarse particles is a concern. For the above reasons, the precipitation amount of AlN after hot rolling is limited to 0.01% or less. The preferred range is 0.005% or less.
なお、ＡｌＮの析出量の化学分析法としては、臭素メタノール溶液で溶解し、０．２μｍのフィルターで残さを採取し、これを化学分析する方法が一般的である。０．２μｍのフィルターを用いても、ろ過の過程で析出物によりフィルターが目詰まりを起こすため、実際には０．２μｍ以下の微細な析出物の抽出も可能である。 In addition, as a chemical analysis method for the precipitation amount of AlN, a method in which a residue is collected with a 0.2 μm filter after being dissolved in a bromine methanol solution, and this is chemically analyzed is generally used. Even when a 0.2 μm filter is used, the filter is clogged by precipitates during the filtration process, so that it is actually possible to extract fine precipitates of 0.2 μm or less.
次に、本発明の請求項２、８では、請求項１、７に加えてＮｂ：０．０４％未満を含有するが、このように限定した理由を以下に述べる。
Ｎｂは浸炭加熱の際に鋼中のＣ、Ｎと結びついてＮｂ（ＣＮ）を形成し、結晶粒の粗大化抑制に有効な元素である。Ｎｂ添加により「Ｔｉ系析出物による粗大粒防止」効果が一層有効になる。これは、Ｔｉ系析出物にＮｂが固溶し、Ｔｉ系析出物の粗大化を抑制するためである。そのため、本願発明の添加量の範囲内では、Ｎｂの添加量に依存して効果は増大するものの、０．０３％未満、あるいは０．０２％未満、さらには０．０１％未満といった微量添加においても、Ｎｂを添加しない場合に比較して、粗大粒防止特性は顕著に向上する。但し、Ｎｂ添加は切削性や冷間鍛造性の劣化、浸炭特性の劣化を引き起こす。特に、Ｎｂの添加量がＮｂ：０．０４％以上であると、素材の硬さが硬くなって切削性、冷間鍛造性が劣化するとともに、棒鋼・線材圧延加熱時の溶体化が困難になる。以上の理由から、その含有量を０．０４％未満の範囲内にする必要がある。切削性、冷間鍛造性等の加工性を重視する場合の好適範囲は、０．０３％未満である。また、加工性に加えて、浸炭性を重視する場合の好適範囲は０．０２％未満である。さらに、特別に浸炭性を重視する場合の好適範囲は０．０１％未満である。なお、Ｎｂの含有量の下限は特に限定するものではないが、０．００１％を下限とすることが好ましい。また、粗大粒防止特性と加工性の両立を図るために、Ｎｂの添加量は、Ｔｉの添加量に応じて、調整することが推奨される。例えば、Ｔｉ＋Ｎｂの好適範囲は、０．０７〜０．１７％未満である。特に高温浸炭や、冷鍛部品において、望ましい範囲は０．０９１％超〜０．１７％未満である。
Next, in claims 2 and 8 of the present invention, in addition to claims 1 and 7 , Nb: less than 0.04% is contained. The reason for this limitation will be described below.
Nb combines with C and N in steel during carburizing heating to form Nb (CN), and is an element effective for suppressing coarsening of crystal grains. By adding Nb, the effect of “preventing coarse grains due to Ti-based precipitates” becomes more effective. This is because Nb dissolves in the Ti-based precipitate and suppresses the coarsening of the Ti-based precipitate. Therefore, within the range of the addition amount of the present invention, the effect increases depending on the addition amount of Nb, but in a small amount addition such as less than 0.03%, less than 0.02%, and even less than 0.01%. However, compared with the case where Nb is not added, the coarse grain prevention characteristic is remarkably improved. However, addition of Nb causes deterioration of machinability, cold forgeability, and carburization characteristics. In particular, when the amount of Nb added is Nb: 0.04% or more, the hardness of the material becomes hard and the machinability and cold forgeability deteriorate, and it is difficult to form a solution during heating of the steel bar and wire rod. Become. For the above reasons, the content needs to be within a range of less than 0.04%. The preferred range when workability such as machinability and cold forgeability is important is less than 0.03%. In addition to workability, the preferred range when carburization is important is less than 0.02%. Furthermore, the preferable range in the case where carburizability is particularly emphasized is less than 0.01%. In addition, although the minimum of content of Nb is not specifically limited, It is preferable to make 0.001% into a minimum. In order to achieve both the coarse grain prevention characteristics and the workability, it is recommended that the amount of Nb added be adjusted according to the amount of Ti added. For example, the preferable range of Ti + Nb is 0.07 to less than 0.17%. Particularly in high-temperature carburizing and cold forged parts, the desirable range is more than 0.091% and less than 0.17%.
次に、本発明の請求項４、１０では、熱間圧延後のベイナイトの組織分率を３０％以下に制限するが、このように限定した理由を以下に述べる。熱間圧延後の鋼材にベイナイト組織が混入すると、浸炭加熱時の粗大粒発生の原因になる。また、ベイナイトの混入の抑制は冷間加工性改善の視点からも望ましい。これらの悪影響は、ベイナイトの組織分率が３０％を超えると特に顕著になる。以上の理由から、熱間圧延後のベイナイトの組織分率を３０％以下に制限する必要がある。高温浸炭等で粗大粒防止に対して浸炭条件が厳しい場合の好適範囲は２０％以下である。冷鍛経由等でさらに粗大粒防止に対して浸炭条件が厳しい場合の好適範囲は１０％以下である。
次に本発明の請求項５、１１では、熱間圧延後のフェライト結晶粒度番号をＪＩＳ Ｇ０５５２で規定されている８〜１１番とするが、このように限定した理由を以下に述べる。熱間圧延後のフェライト粒が過度に微細であると、浸炭時にオーステナイト粒が過度に微細化する。オーステナイト粒が過度に微細になると、粗大粒が生成しやすくなり、特にフェライト結晶粒度が１１番を超えるとその傾向が顕著になる。また、オーステナイト粒がＪＩＳ Ｇ０５５１で規定されている１１番を超えるような過度に微細になると、前掲の特開平２００３−３４８４３公報の鋼材と同様に、焼入れ性の劣化による強度不足等の弊害を生じる。一方、熱間圧延後のフェライト結晶粒度番号をＪＩＳ Ｇ０５５２で規定されている８番未満の粗粒にすると、熱間圧延材の延性が劣化し、冷間鍛造性が劣化する。以上の理由から、熱間圧延後のフェライト結晶粒度番号をＪＩＳ Ｇ０５５２で規定されている８〜１１番の範囲内にする必要がある。
Next, in claims 4 and 10 of the present invention, the structure fraction of bainite after hot rolling is limited to 30% or less. The reason for this limitation will be described below. If a bainite structure is mixed in the steel after hot rolling, coarse grains are generated during carburizing heating. In addition, suppression of bainite contamination is also desirable from the viewpoint of improving cold workability. These adverse effects become particularly prominent when the bainite structural fraction exceeds 30%. For the above reasons, it is necessary to limit the structure fraction of bainite after hot rolling to 30% or less. The preferred range when the carburizing conditions are severe for preventing coarse grains due to high-temperature carburizing or the like is 20% or less. The preferred range is 10% or less when the carburizing conditions are severe for preventing coarse grains via cold forging.
Next , in claims 5 and 11 of the present invention , the ferrite grain size number after hot rolling is set to Nos. 8 to 11 defined in JIS G0552. The reason for this limitation will be described below. If the ferrite grains after hot rolling are excessively fine, the austenite grains are excessively refined during carburizing. If the austenite grains become excessively fine, coarse grains are likely to be formed, and the tendency becomes prominent particularly when the ferrite crystal grain size exceeds # 11. In addition, when the austenite grains become excessively fine such as exceeding No. 11 defined in JIS G0551, problems such as insufficient strength due to deterioration of hardenability occur as in the steel material disclosed in JP-A-2003-34843. . On the other hand, when the ferrite crystal grain size number after hot rolling is coarser than the number 8 specified in JIS G0552, the ductility of the hot rolled material is deteriorated and the cold forgeability is deteriorated. For the above reasons, it is necessary to set the ferrite crystal grain size number after hot rolling within the range of 8 to 11 defined in JIS G0552.
本発明の請求項６では、熱間圧延後の鋼のマトリックス中の長手方向断面において、検査基準面積：１００平方ｍｍ、検査回数１６視野、予測を行なう面積：３００００平方ｍｍの条件で測定された極値統計によるＴｉ系析出物の最大直径が４０μｍ以下とするが、このように限定した理由を以下に述べる。本発明で対象とする浸炭部品の要求特性の一つに転動疲労特性や面疲労強度のような接触疲労強度がある。粗大なＴｉ系析出物が鋼中に存在すると接触疲労破壊の起点となり、疲労特性が劣化する。極値統計により、検査基準面積：１００平方ｍｍ、検査回数１６視野、予測を行なう面積：３００００平方ｍｍの条件で測定した時のＴｉ系析出物の最大直径が４０μｍを超えると、特に接触疲労特性に及ぼすＴｉ系析出物の悪影響が顕著になる。以上の理由から、極値統計によるＴｉ系析出物の最大直径を４０μｍ以下とする必要がある。極値統計による析出物の最大直径の測定・予測方法は、１９９３年３月８日養賢堂発行「金属疲労 微小欠陥と介在物の影響」２３３頁〜２５２頁に記載の方法による。なお、本発明で用いているのは、二次元的検査により一定面積内（予測を行なう面積：３００００平方ｍｍ）で観察される最大析出物を推定するという二次元的検査方法である。詳細な測定手順は、実施例欄で述べる。
According to claim 6 of the present invention, in the longitudinal section in the matrix of the steel after hot rolling, it was measured under the conditions of inspection standard area: 100 square mm, inspection number of view 16 fields, prediction area: 30000 square mm. The maximum diameter of the Ti-based precipitate by extreme value statistics is set to 40 μm or less. The reason for this limitation will be described below. One of the required characteristics of the carburized parts targeted by the present invention is contact fatigue strength such as rolling fatigue characteristics and surface fatigue strength. If coarse Ti-based precipitates are present in the steel, it becomes a starting point for contact fatigue failure, and the fatigue characteristics deteriorate. According to extreme value statistics, when the maximum diameter of Ti-based precipitates exceeds 40 μm when measured under the conditions of inspection standard area: 100 square mm, number of inspections of 16 fields, and prediction area: 30000 square mm, especially contact fatigue characteristics The adverse effect of Ti-based precipitates on the surface becomes remarkable. For the above reasons, the maximum diameter of Ti-based precipitates according to extreme value statistics needs to be 40 μm or less. The method of measuring and predicting the maximum diameter of precipitates by extreme value statistics is based on the method described on pages 233 to 252 of “Effects of Metal Fatigue Microdefects and Inclusions” issued by Yokendo on March 8, 1993. In addition, what is used in the present invention is a two-dimensional inspection method in which the maximum precipitate observed within a certain area (predicted area: 30000 square mm) is estimated by a two-dimensional inspection. Detailed measurement procedures are described in the Examples section.
次に熱間圧延条件について説明する。
上記の本発明成分からなる鋼を、転炉、電気炉等の通常の方法によって溶製し、成分調整を行い、鋳造工程、必要に応じて分塊圧延工程を経て、線材または棒鋼に熱間圧延する圧延素材とする。
Next, hot rolling conditions will be described.
The steel composed of the above-described components of the present invention is melted by a usual method such as a converter, an electric furnace, etc., and the components are adjusted. A rolling material to be rolled is used.
次に、本発明の請求項７は、圧延素材を１１５０℃以上の温度で保熱時間１０分以上加熱の温度で加熱する。加熱条件が、１１５０℃未満であるか、あるいは加熱温度が１１５０℃以上でも保熱時間が１０分未満では、Ｔｉ系析出物、Ｎｂの析出物およびＡｌＮを一旦マトリックス中に固溶させることができない。そのため、熱間圧延後の鋼材に、一定量以上のＴｉ系析出物、Ｎｂの析出物をあらかじめ微細析出させることができず、熱間圧延後に粗大なＴｉ系析出物、Ｎｂの析出物、ＡｌＮが存在し、浸炭時に粗大粒の発生を抑制することができない。そのため、熱間圧延に際して、１１５０℃以上の温度で保熱時間１０分以上加熱することが必要である。好適範囲は１１８０℃以上の温度で保熱時間１０分以上である。
Next, Claim 7 of this invention heats a rolling raw material at the temperature of 1150 degreeC or more and the heat retention time for 10 minutes or more. If the heating condition is less than 1150 ° C. or the heat retention time is less than 10 minutes even if the heating temperature is 1150 ° C. or higher, Ti-based precipitates, Nb precipitates and AlN cannot be once dissolved in the matrix. . Therefore, a certain amount of Ti-based precipitates and Nb precipitates cannot be finely precipitated in advance on the steel material after hot rolling, and coarse Ti-based precipitates, Nb precipitates, and AlN after hot rolling. And the generation of coarse particles during carburization cannot be suppressed. Therefore, in hot rolling, it is necessary to heat at a temperature of 1150 ° C. or higher for a heat retention time of 10 minutes or longer. A preferable range is a temperature of 1180 ° C. or higher and a heat retention time of 10 minutes or longer.
次に、本発明の請求項１０は、熱間圧延後に８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する。冷却条件は、１℃／ｓを越えるとＴｉ系析出物の析出温度域を短時間しか通過させることができず、熱間圧延後の微細なＴｉＣ系析出物の析出量が不十分となり、さらにまた、ベイナイトの組織分率が大きくなる。そのため、浸炭時に粗大粒の発生を抑制することができなくなる。また、冷却速度が大きいと圧延材の硬さが上昇し、冷間鍛造性が劣化するため、冷却速度はできるだけ小さくするのが望ましい。好適範囲は０．７℃／ｓ以下である。なお、冷却速度を小さくする方法としては、圧延ラインの後方に保温カバーまたは熱源付き保温カバーを設置し、これにより、徐冷を行う方法が挙げられる。
Next, Claim 10 of this invention anneals the temperature range of 800-500 degreeC with the cooling rate of 1 degrees C / sec or less after hot rolling. When the cooling condition exceeds 1 ° C./s, the precipitation temperature region of the Ti-based precipitate can be passed only for a short time, and the amount of fine TiC-based precipitate after hot rolling becomes insufficient, Moreover, the structure fraction of bainite increases. Therefore, it becomes impossible to suppress the generation of coarse particles during carburizing. Further, if the cooling rate is high, the hardness of the rolled material increases and the cold forgeability deteriorates. Therefore, it is desirable to make the cooling rate as small as possible. The preferred range is 0.7 ° C./s or less. In addition, as a method of reducing the cooling rate, a method of installing a heat insulating cover or a heat insulating cover with a heat source behind the rolling line and thereby performing slow cooling can be mentioned.
次に、本発明の請求項１１は、熱間圧延の仕上げ温度を８４０〜１０００℃とする。仕上げ温度が８４０℃未満では、フェライト結晶粒度が過度に微細になりすぎて、その後の浸炭時に粗大粒が発生しやすくなる。一方、仕上げ温度が１０００℃を超えると、圧延材の硬さが硬くなって冷間鍛造性が劣化する。以上の理由から、熱間圧延の仕上げ温度を８４０〜１０００℃とする。冷間鍛造用途で、焼鈍なしで使用する場合は、８４０〜９２０℃の範囲が、それ以外では９２０〜１０００℃の範囲が望ましい。 Next, claim 11 of the present invention sets the finishing temperature of hot rolling to 840 to 1000 ° C. If the finishing temperature is less than 840 ° C., the ferrite crystal grain size becomes excessively fine, and coarse grains are likely to be generated during the subsequent carburizing. On the other hand, when the finishing temperature exceeds 1000 ° C., the hardness of the rolled material becomes hard and the cold forgeability deteriorates. For these reasons, the hot rolling finishing temperature is set to 840 to 1000 ° C. When used for cold forging without annealing, a range of 840 to 920 ° C is desirable, and a range of 920 to 1000 ° C is desirable otherwise.
本発明では、鋳片のサイズ、凝固時の冷却速度、分塊圧延条件については特に限定するものではなく、本発明の要件を満足すればいずれの条件でも良い。
本発明は、冷間鍛造工程で製造される部品、熱間鍛造で製造される部品いずれにも適用可能である。熱間鍛造工程の例としては、「棒鋼−熱間鍛造−必要により焼準等の熱処理−切削−浸炭焼入れ−必要により研磨」の工程があげられる。本発明の鋼材を用いて、例えば１１５０℃以上の加熱温度で熱間鍛造を行い、その後必要に応じて焼準処理を行なうことにより、９５０℃〜１０９０℃の温度域での浸炭のような厳しい条件での浸炭焼入れ熱処理においても、粗大粒の発生の抑制が可能となり、優れた材質特性が得られる。例えば、軸受部品、転動部品の場合であると、高温浸炭を行っても、優れた転動疲労特性が得られる。
本発明では、浸炭条件を特に限定するものではない。軸受部品、転動部品において、特に高いレベルの転動疲労寿命を指向する場合には、上記のように、浸炭時の炭素ポテンシャルを０．９〜１．３％の範囲で高めに設定すること、あるいは、いわゆる浸炭浸窒処理を行うことが有効である。浸炭浸窒処理は、浸炭後の拡散処理の過程で浸窒を行う処理であるが、表面の窒素濃度が０．２〜０．６％の範囲になるような条件が適切である。これらの条件を選択することにより、浸炭層に微細なＴｉ（ＣＮ）が多量に析出し、また残留γが３０〜４０％導入されることが、転動寿命の向上に寄与する。
In the present invention, the size of the slab, the cooling rate during solidification, and the ingot rolling conditions are not particularly limited, and any conditions may be used as long as the requirements of the present invention are satisfied.
The present invention can be applied to both parts manufactured by the cold forging process and parts manufactured by hot forging. As an example of the hot forging process, there is a process of “bar steel-hot forging-heat treatment such as normalization if necessary-cutting-carburizing and quenching-polishing if necessary". By using the steel material of the present invention, for example, hot forging at a heating temperature of 1150 ° C. or higher, and then performing a normalization treatment as necessary, such as severe carburization in a temperature range of 950 ° C. to 1090 ° C. Even in carburizing and quenching heat treatment under conditions, the generation of coarse grains can be suppressed, and excellent material properties can be obtained. For example, in the case of bearing parts and rolling parts, excellent rolling fatigue characteristics can be obtained even when high-temperature carburizing is performed.
In the present invention, carburizing conditions are not particularly limited. In bearing parts and rolling parts, when aiming at a particularly high level of rolling fatigue life, as described above, the carbon potential during carburizing should be set higher within the range of 0.9 to 1.3%. Alternatively, it is effective to perform a so-called carburizing and nitriding treatment. The carburizing and nitriding treatment is a treatment in which nitriding is performed in the course of the diffusion treatment after carburizing, and conditions under which the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. By selecting these conditions, a large amount of fine Ti (CN) precipitates in the carburized layer, and the introduction of 30 to 40% of residual γ contributes to the improvement of the rolling life.
以下に、本発明の効果を実施例により、さらに具体的に示す。 Hereinafter, the effects of the present invention will be described more specifically by way of examples.
表１に示す組成を有する転炉溶製鋼を連続鋳造し、必要に応じて分塊圧延工程を経て１６２ｍｍ角の圧延素材とした。続いて、熱間圧延により、直径２４〜３０ｍｍの棒鋼を製造した。 Converter molten steel having the composition shown in Table 1 was continuously cast, and a rolling raw material of 162 mm square was obtained through a batch rolling process as necessary. Subsequently, steel bars having a diameter of 24 to 30 mm were manufactured by hot rolling.
熱間圧延後の棒鋼から、ＡｌＮの析出量を化学分析により求めた。また、圧延後の棒鋼について、ミクロ観察を行い、ベイナイト分率の測定、ＪＩＳ Ｇ ０５５２の規定に従ってフェライト結晶粒度の測定を行なった。さらに、ビッカース硬さを測定し、冷間加工性の指標とした。 From the steel bar after hot rolling, the precipitation amount of AlN was determined by chemical analysis. Further, the rolled steel bar was micro-observed, the bainite fraction was measured, and the ferrite crystal grain size was measured in accordance with JIS G 0552. Furthermore, the Vickers hardness was measured and used as an index of cold workability.
上記の工程で製造した棒鋼について、球状化焼鈍を行った後、据え込み試験片を作成し、圧下率５０％の据え込みを行った後、浸炭シミュレーションを行った。浸炭シミュレーションの条件は、９１０℃〜１０１０℃に５時間加熱−水冷である。その後、切断面に研磨−腐食を行い、旧オーステナイト粒径を観察して粗粒発生温度（結晶粒粗大化温度）を求めた。浸炭処理は通常９３０〜９５０℃の温度域で行われるため、粗粒発生温度が９５０℃以下のものは結晶粒粗大化特性に劣ると判定した。なお、旧オーステナイト粒度の測定はＪＩＳ Ｇ ０５５１に準じて行い、４００倍で１０視野程度観察し、粒度番号５番以下の粗粒が１つでも存在すれば粗粒発生と判定した。
また、極値統計法によるＴｉ系析出物の最大直径の予測は次の方法で行なった。析出物がＴｉ系であるか否かは、光学顕微鏡におけるコントラストの違いからを判別した。コントラストの違いによる判別法の妥当性は、あらかじめエネルギー分散型Ｘ線分光分析装置付き走査型電子顕微鏡にて確認した。試験片長手方向断面において検査基準面積１００平方ｍｍ（１０ｍｍ×１０ｍｍの領域）の領域をあらかじめ１６視野分準備した。そして各検査基準面積１００平方ｍｍにおけるＴｉ系の最大析出物を検出し、これを光学顕微鏡にて１０００倍で写真撮影した。これを各々の各検査基準面積１００平方ｍｍの１６視野について、１６回繰り返し行なった（つまり検査回数１６視野）。得られた写真から各検査基準面積における最大析出物の直径を計測した。楕円形の場合は長径と短径の相乗平均を求めその析出物の直径とした。得られた最大析出物直径の１６個のデータを、前記の養賢堂発行「金属疲労微小欠陥と介在物の影響」２３３頁〜２５２頁に記載の方法に従い、最大析出物分布直線（最大析出物直径と極値統計基準化変数の一次関数）を最小二乗法により求め、最大析出物分布直線を外挿することにより、予測を行なう面積：３００００平方ｍｍにおける最大析出物の直径を予測した。
The steel bar manufactured in the above process was subjected to spheroidizing annealing, then an upsetting test piece was prepared, and upsetting was performed at a reduction rate of 50%, and then a carburizing simulation was performed. The conditions for the carburizing simulation are heating to 910 ° C. to 1010 ° C. for 5 hours and water cooling. Thereafter, the cut surface was polished and corroded, and the prior austenite grain size was observed to determine the coarse grain generation temperature (crystal grain coarsening temperature). Since the carburizing process is normally performed in a temperature range of 930 to 950 ° C., it was determined that the coarse grain generation temperature was 950 ° C. or lower and the crystal grain coarsening characteristics were inferior. The prior austenite particle size was measured in accordance with JIS G 0551, observed at 400 times for about 10 fields of view, and if there was at least one coarse particle having a particle size number of 5 or less, it was determined that coarse particles were generated.
Moreover, the maximum diameter of Ti-based precipitates was estimated by the following method using the extreme value statistical method. Whether or not the precipitate is Ti-based was determined from the difference in contrast in the optical microscope. The validity of the discrimination method based on the difference in contrast was confirmed in advance with a scanning electron microscope equipped with an energy dispersive X-ray spectrometer. An area having an inspection reference area of 100 square mm (10 mm × 10 mm area) in the longitudinal section of the test piece was prepared for 16 visual fields in advance. And the Ti-type largest deposit in each inspection reference area 100 square mm was detected, and this was photographed 1000 times with the optical microscope. This was repeated 16 times for each of the 16 visual fields of each inspection reference area 100 square mm (that is, the number of inspections was 16 visual fields). The diameter of the largest deposit in each inspection reference area was measured from the obtained photograph. In the case of an ellipse, the geometric mean of the major axis and minor axis was determined and used as the diameter of the precipitate. According to the method described in “Effects of metal fatigue microdefects and inclusions” on pages 233 to 252 published by Yokendo, the maximum data of the maximum precipitate diameter was obtained. A linear function of the object diameter and the extreme value statistical normalization variable) was obtained by the least square method, and the maximum precipitate distribution line was extrapolated from the maximum precipitate distribution line, thereby predicting the maximum precipitate diameter in the area to be predicted: 30000 square mm.
次に、圧下率５０％で冷間鍛造を行なった各鋼材から、直径１２．２ｍｍの円柱状の転動疲労試験片を作成し、９５０℃×５時間、炭素ポテンシャル０．８％の条件で浸炭を行なった。焼入れ油の温度は１３０℃、焼戻しは１８０℃×２時間である。これらの浸炭焼入れ材について、浸炭層のγ粒度を調査した。さらに、点接触型転動疲労試験機（ヘルツ最大接触応力５８８４ＭＰａ）を用いて転動疲労特性を評価した。疲労寿命の尺度として、「試験結果をワイブル確率紙にプロットして得られる累積破損確率１０％における疲労破壊までの応力繰り返し数」として定義されるＬ10寿命を用いた。
これらの調査結果をまとめて、表２に示す。転動疲労寿命は比較例１２のＬ10寿命を１とした時の各材料のＬ10寿命の相対値を示した。
本発明例の結晶粒粗大化温度は９９０℃以上であり、９５０℃浸炭材のγ粒も細整粒であり、転動疲労特性もすぐれていることが明らかである。
一方、比較例１２,１３はＪＩＳのＳＣｒ４２０およびＳＣＭ４２０の特性であるが粗大粒発生温度は低く、９５０℃浸炭材のγ粒が粗大化している。また、比較例１１はＮの含有量が本願規定の範囲を上回りさらにＴｉ系析出物の最大直径が本願規定の範囲を上回った場合であり、粗大粒の生成が顕著に見られるとともに、転動疲労特性も良くない。比較例１４はＴｉ含有量が本願規定の範囲を下回った場合であり、Ｔｉのピン止め効果が小さいため、粗大粒の抑制に効果を表していない。比較例１５はＴｉ含有量が本願規定の範囲を上回り、さらにＴｉ系析出物の最大直径が本願規定の範囲を上回った場合であり、ＴｉＣによる析出効果が顕著に見られ、冷間加工性の劣化を招く。また、Ｔｉ系析出物の溶体化不良を招くために、粗大粒防止特性も劣り、転動疲労特性も良くない。比較例１６はＮｂ含有量が本願規定の範囲を上回った場合であり、素材の硬さが硬くなり冷間加工性の劣化を招くとともに、粗大粒防止特性も劣る。比較例１７はＯ含有量が本願規定の範囲を上回った場合であり、これも粗大粒が生成し、転動疲労特性も良くない。比較例１８は成分は本願規定の範囲内であるが、熱間圧延後の冷却速度が本発明の範囲を上回り、熱間圧延後のベイナイト組織分率が本願規定の範囲を超えており、これも粗大粒が生成する。比較例１９は仕上げ温度が本発明の範囲を上回り、圧延後のフェライト結晶粒度が本発明の範囲より粗大となった場合であり、粗大粒防止特性は劣る。比較例２０は仕上げ温度が本発明の範囲を下回り、圧延後のフェライト結晶粒度が本発明の範囲より微細となった場合であり、やはり粗大粒防止特性は劣る。比較例２１は圧延加熱温度が本発明の範囲を下回り、熱間圧延後のＡｌＮの析出量が本願規定の範囲を上回った場合であり、粗大粒防止特性は劣り、転動疲労特性も良くない。
Next, a cylindrical rolling fatigue test piece having a diameter of 12.2 mm was prepared from each steel material that had been cold forged at a reduction rate of 50%, and was subjected to conditions of 950 ° C. × 5 hours and a carbon potential of 0.8%. Carburized. The temperature of the quenching oil is 130 ° C., and the tempering is 180 ° C. × 2 hours. For these carburized and quenched materials, the γ grain size of the carburized layer was investigated. Furthermore, the rolling fatigue characteristics were evaluated using a point contact type rolling fatigue tester (Hertz maximum contact stress 5884 MPa). As a measure of fatigue life, L 10 life defined as “the number of stress repetitions until fatigue failure at a cumulative failure probability of 10% obtained by plotting test results on Weibull probability paper” was used.
These survey results are summarized in Table 2. The rolling fatigue life indicates the relative value of the L 10 life of each material when the L 10 life of Comparative Example 12 is 1.
It is clear that the crystal grain coarsening temperature of the example of the present invention is 990 ° C. or higher, the γ grains of the 950 ° C. carburized material are finely sized, and have excellent rolling fatigue characteristics.
On the other hand, Comparative Examples 12 and 13 are the characteristics of JIS SCr420 and SCM420, but the coarse grain generation temperature is low, and the γ grains of the 950 ° C. carburized material are coarsened. Further, Comparative Example 11 is a case where the N content exceeds the range specified in the present application and the maximum diameter of the Ti-based precipitate exceeds the range specified in the present application. Fatigue properties are not good. Comparative example 14 is a case where the Ti content falls below the range specified in the present application, and since the pinning effect of Ti is small, it does not show an effect in suppressing coarse grains. Comparative Example 15 is a case where the Ti content exceeds the range specified in the present application, and the maximum diameter of the Ti-based precipitate exceeds the range specified in the present application. The precipitation effect due to TiC is noticeable, and the cold workability is low. It causes deterioration. Moreover, since it causes poor solution of Ti-based precipitates, the coarse grain prevention characteristics are inferior and the rolling fatigue characteristics are not good. Comparative Example 16 is a case where the Nb content exceeds the range specified in the present application, and the hardness of the material becomes hard, resulting in deterioration of cold workability and inferior coarse grain prevention characteristics. Comparative Example 17 is a case where the O content exceeds the range specified in the present application, and this also produces coarse grains and poor rolling fatigue characteristics. In Comparative Example 18, the components are within the range specified in the present application, but the cooling rate after hot rolling exceeds the range of the present invention, and the bainite structure fraction after hot rolling exceeds the range specified in the present application. Coarse grains are also produced. Comparative Example 19 is a case where the finishing temperature exceeds the range of the present invention, and the ferrite crystal grain size after rolling becomes coarser than the range of the present invention, and the coarse grain preventing property is inferior. Comparative Example 20 is a case where the finishing temperature falls below the range of the present invention, and the ferrite crystal grain size after rolling becomes finer than the range of the present invention, and the coarse grain prevention property is also inferior. Comparative Example 21 is a case where the rolling heating temperature is lower than the range of the present invention, and the precipitation amount of AlN after hot rolling exceeds the range specified in the present application, the coarse grain prevention characteristics are inferior, and the rolling fatigue characteristics are not good. .
表３に示す組成を有する転炉溶製鋼を連続鋳造し、必要に応じて分塊圧延工程を経て１６２ｍｍ角の圧延素材とした。続いて、熱間圧延により、直径７０ｍｍの棒鋼を製造した。この棒鋼を素材として、熱間鍛造を行い直径４０ｍｍの熱間鍛造部材に仕上げた。熱間鍛造の加熱温度は１１００℃〜１２９０℃である。
上記の工程で製造した熱間鍛造部材について、９００℃×１時間加熱空冷の条件で焼準処理を行った。その後、加熱時間５時間の条件で浸炭シミュレーションを行い、実施例−１と同様に、粗大粒発生温度を求めた。
また、上記の熱間鍛造部材を焼準した後、直径１２．２ｍｍの円柱状の転動疲労試験片を作成し、１０５０℃×１時間、炭素ポテンシャル１．０％の条件で浸炭焼入れを行った。焼入れ油の温度は１３０℃、焼戻しは１８０℃×２時間の条件である。
これらの調査結果をまとめて、表４に示す。転動疲労寿命は比較例１２のＬ10寿命を１とした時の各材料のＬ10寿命の相対値を示した。
表４に示した通り、本発明例では、結晶粒粗大化温度は１０７０℃超である。また、１０５０℃浸炭材のγ粒は８番以上の細粒であり、転動疲労寿命も比較例に比べて２倍以上と極めて良好である。
一方、比較例は、実施例−１と同様に本発明の要件の範囲から逸脱しており、粗大粒防止特性は劣り、転動疲労特性も良くない。
Converter molten steel having the composition shown in Table 3 was continuously cast, and a rolling raw material of 162 mm square was obtained through a batch rolling process as necessary. Subsequently, a steel bar having a diameter of 70 mm was manufactured by hot rolling. Using this steel bar as a raw material, hot forging was performed to finish a hot forged member having a diameter of 40 mm. The heating temperature of hot forging is 1100 ° C to 1290 ° C.
About the hot forging member manufactured at said process, the normalization process was performed on the conditions of 900 degreeC x 1 hour heating and air cooling. Then, carburizing simulation was performed on the conditions of heating time 5 hours, and the coarse grain generation temperature was calculated | required similarly to Example-1.
In addition, after normalizing the above hot forged member, a cylindrical rolling fatigue test piece having a diameter of 12.2 mm was prepared and carburized and quenched at 1050 ° C. for 1 hour with a carbon potential of 1.0%. It was. The temperature of the quenching oil is 130 ° C., and the tempering is 180 ° C. × 2 hours.
These survey results are summarized in Table 4. The rolling fatigue life indicates the relative value of the L 10 life of each material when the L 10 life of Comparative Example 12 is 1.
As shown in Table 4, in the present invention example, the crystal grain coarsening temperature is over 1070 ° C. Further, the γ grains of the 1050 ° C. carburized material are fine grains of No. 8 or more, and the rolling fatigue life is extremely good, that is, twice or more as compared with the comparative example.
On the other hand, the comparative example deviates from the range of the requirements of the present invention as in Example-1 and the coarse grain prevention property is inferior and the rolling fatigue property is not good.
Claims (11)
Ｃ：０．１〜０．４％、
Ｓｉ：０．０２〜１．５％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．００５〜０．０５％
Ｔｉ：０．０５〜０．２％、
Ｎ：０．００５１％未満に制限し、
Ｃｒ：０．４〜２．０％、
Ｐ：０．０２５％以下に制限し、
Ｏ：０．００２５％以下に制限し、
残部が鉄および不可避的不純物からなり、
熱間圧延後のＡｌＮの析出量を０．０１％以下に制限したことを特徴とする浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼。 As mass%
C: 0.1-0.4%
Si: 0.02 to 1.5%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.005 to 0.05%
Ti: 0.05-0.2 %
N : limited to less than 0.0051% ,
Cr : 0.4 to 2.0 %,
P : limited to 0.025% or less,
O: limited to 0.0025% or less,
The balance consists of iron and inevitable impurities,
A case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing, characterized by limiting the precipitation amount of AlN after hot rolling to 0.01% or less.
Ｃ：０．１〜０．４％、
Ｓｉ：０．０２〜１．５％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．００５〜０．０５％
Ｔｉ：０．０５〜０．２％、
Ｎ：０．００５１％未満に制限し、
Ｃｒ：０．４〜２．０％、
Ｐ：０．０２５％以下に制限し、
Ｏ：０．００２５％以下に制限し、
残部が鉄および不可避的不純物からなる鋼を、
１１５０℃以上の温度で保熱時間１０分以上加熱して線材または棒鋼に熱間圧延し、熱間圧延後のＡｌＮの析出量を０．０１％以下に制限した鋼となるようにすることを特徴とする浸炭時の粗大粒防止特性と疲労特性に優れた肌焼鋼の製造方法。 As mass%
C: 0.1-0.4%
Si: 0.02 to 1.5%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.005 to 0.05%
Ti: 0.05-0.2 %
N : limited to less than 0.0051% ,
Cr : 0.4 to 2.0 %,
P : limited to 0.025% or less,
O: limited to 0.0025% or less,
Steel with the balance of iron and inevitable impurities
Heating at a temperature of 1150 ° C. or higher for 10 minutes or longer and hot rolling to a wire or steel bar, so that the precipitation amount of AlN after hot rolling is limited to 0.01% or less. A method for producing case-hardened steel with excellent coarse grain prevention characteristics and fatigue characteristics during carburizing.
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