JP2004183064A - Steel for case hardening having excellent cold workability and coarse grain prevention property when carburized, and production method therefor - Google Patents
Steel for case hardening having excellent cold workability and coarse grain prevention property when carburized, and production method therefor Download PDFInfo
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【０００１】
【発明の属する技術分野】
本発明は、冷間加工性に優れ、かつ浸炭時の粗大粒防止特性に優れた肌焼用鋼材およびその製造方法に関する。
【０００２】
【従来の技術】
歯車、シャフト、ＣＶＪ部品は、通常、例えばＪＩＳ Ｇ ４０５２、ＪＩＳＧ 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）−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。ここで、熱間鍛造から冷間鍛造への切り替えに際しては、鋼材の冷間変形抵抗の低減と限界圧縮率の向上が重要な課題である。これは、前者は、鍛造工具の寿命を確保するためであり、後者は冷間鍛造時の鋼材の割れを防止するためである。このような冷間鍛造に適した鋼材として、ボロン鋼の適用が検討されている。しかしながら、ボロン鋼は浸炭加熱時に一部のオーステナイト結晶粒が粗大化する現象を起こしやすい。このような粗大粒が発生した部品では、浸炭焼入れ後に熱処理歪みを発生し、例えば、歯車やＣＶＪ部品ではこの浸炭歪みが大きければ、騒音や振動の原因となる。こうした経緯から、冷間加工性に優れ、かつ低浸炭歪み特性に優れた、つまり浸炭時に粗大粒を生じないボロン鋼が強く求められている。
【０００３】
本出願人は、先に冷間加工性に優れ、かつ浸炭時に粗大粒の発生を防止できて低浸炭歪み特性に優れた肌焼ボロン鋼とその製造方法の発明を提供した（特許文献１）。
【０００４】
この発明では、鋼の化学成分を制限するのみでは、粗大粒を防止できないので、熱間加工後の鋼材のＡｌＮの固溶状態や、Ｎｂ（ＣＮ）、Ｔｉ（ＣＮ）等の微細析出物の析出量を制御して、ＮｂＣ、ＴｉＣ析出物をピン止め粒子として活用する手段等を採用している。
【０００５】
即ち、圧延加熱温度を１１５０℃以上の高温にして、熱間加工後の鋼材の状態でＡｌＮの析出量を制限し、また、圧延加熱温度を１１５０℃以上の高温とすることによりＮｂの析出物、Ｔｉの析出物を一旦溶体化し、熱間圧延後にこれらの析出物の析出温度域を徐冷することにより、熱間圧延後の鋼材に一定量以上のＮｂの析出物、Ｔｉの析出物を多量微細分散させるものである。
【０００６】
また、本出願人は、冷間鍛造時には冷間加工性に優れ、浸炭時に粗大粒の発生と表面から深さ０．２〜０．７ｍｍに生成する不完全焼入れ組織の生成を防止することができる肌焼ボロン鋼とその製造方法の発明を提供した（特許文献２）。
【０００７】
この発明は、化学組成を制限した鋼を加熱温度を高温の１１５０℃以上とし、熱間圧延の仕上温度を８４０〜１０００℃、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷することを特徴としている。
【０００８】
上記いずれの発明も、圧延加熱温度を１１５０℃以上の高温としてＴｉの析出物、Ｎｂの析出物を一旦溶体化し、熱間圧延後に析出温度域を徐冷することにより、これらの析出物を多量微細分散させているものである。しかしながら、これらの方法でも浸炭時の粗大粒の防止が必ずしも十分でない場合があり、さらに、圧延加熱温度を１１５０℃以上の高温にすると、圧延ままで硬くなり、冷間加工性が不十分となり、また、全脱炭が顕著になるといった問題点を有している。
【０００９】
【特許文献１】
特開平１１−３３５７７７号公報
【特許文献２】
特開２００１−３０３１７２号公報
【発明が解決しようとする課題】
本発明は、冷間加工性を向上させ、かつ浸炭時の粗大粒の発生を防止することができる肌焼用鋼材及びその製造方法を提供することを課題とするものである。
【００１０】
【課題を解決しようとする手段】
本発明者は、肌焼用ボロン鋼の冷間加工性の向上及び浸炭時の粗大粒防止特性の向上を図るべく鋭意研究を進めた。その結果、従来のように圧延加熱温度を１１５０℃以上の高温として、全脱炭が顕著になると、その後の浸炭時に脱炭部から粗大粒が発生するため、粗大粒防止特性が不十分となることを見出した。さらに、従来のように圧延加熱温度を１１５０℃以上の高温として、Ｔｉの析出物、Ｎｂの析出物を一旦溶体化しなくても、鋳造後、Ａ３点温度以下に冷却することなくＨＣＲで分塊圧延をした鋼片を低温加熱圧延すれば、全脱炭の低減をすることができ、これにより粗大粒の生成を防止でき、かつ、析出強化量の低減による軟質化により冷間加工性が向上できることを見出して、本発明を完成した。
【００１１】
本発明の要旨は、次のとおりである。
【００１２】
（１） 質量％で、
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなり、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
【００１３】
（２） 質量％で、
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎｂ：０．００２〜０．０５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなり、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＮｂの析出物、Ｔｉの析出物、またはＮｂとＴｉの複合組成からなる析出物のその合計で１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
【００１４】
（３） さらに、質量％で、
Ｍｏ：０．３％以下、
Ｎｉ：２．５％以下
のうちの１種または２種を含有することを特徴とする上記（１）または（２）記載の冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。
【００１５】
（４） 質量％で、
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなる鋼を鋳造後Ａ３点温度以下に冷却することなく分塊圧延を行う工程により製造された鋼片を用い、加熱温度を９００〜１０７０℃、熱間圧延の仕上温度を８００〜９７０℃、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する条件により線材または棒鋼に熱間圧延し、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上とし、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
【００１６】
（５） 質量％で、
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎｂ：０．００２〜０．０５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなる鋼を鋳造後Ａ３点温度以下に冷却することなく分塊圧延を行う工程により製造された鋼片を用い、加熱温度を９００〜１０７０℃、熱間圧延の仕上温度を８００〜９７０℃、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する条件により線材または棒鋼に熱間圧延し、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＮｂの析出物、Ｔｉの析出物、またはＮｂとＴｉの複合組成からなる析出物のその合計で１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
【００１７】
（６） さらに、質量％で、
Ｍｏ：０．３％以下、
Ｎｉ：２．５％以下
のうちの１種または２種を含有することを特徴とする上記（４）または（５）記載の冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。
【００１８】
【発明の実施の形態】
以下、本発明について詳細に説明する。
【００１９】
本発明では、ＨＣＲ（Ｈｏｔ Ｃｈａｒｇｅ Ｒｏｌｌｉｎｇ）により鋼を鋳造後Ａ３点温度以下に冷却することなく分塊圧延を行って製造した鋼片を用い、棒鋼線材圧延に際して低温加熱圧延すると、ＴｉＣの微細析出物がそのまま析出した状態で保持される。析出強化量は、高温で固溶体から析出された析出強化量よりも低減する。また、全脱炭量も低減し、ＪＩＳ Ｇ ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下とすることができる。これによって粗大粒の発生を防止できると共に軟質化により冷間加工性が向上できる。
【００２０】
まず、成分の限定理由について説明する。
【００２１】
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、０．３％を超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、０．１〜０．３％の範囲内にする必要がある。
【００２２】
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、０．０１％未満ではその効果は不十分である。一方、０．１５％を超えると、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０．０１〜０．１５％の範囲内にする必要がある。好適範囲は０．０２〜０．１％である。
【００２３】
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、０．２％未満では効果は不十分であり、０．６５％を超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、０．２〜０．６５％の範囲内にする必要がある。好適範囲は０．３〜０．６％である。
【００２４】
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、０．００５％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０．００５〜０．１５％の範囲内にする必要がある。好適範囲は０．００５〜０．０４％である。
【００２５】
Ｃｒは鋼に強度、焼入れ性を与えるのに有効な元素であるが、０．４％未満ではその効果は不十分であり、１．２５％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０．４〜１．２５％の範囲内にする必要がある。好適範囲は０．６〜１．０％である。
【００２６】
Ｂは次の３点を狙いとして添加する。▲１▼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、圧延ままでの軟質化を促進する。▲２▼浸炭焼入れに際して、鋼に焼入れ性を付与する。▲３▼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。０．０００５％未満の添加では、上記の効果は不十分であり、０．００５％を超えるとその効果は飽和するので、その含有量を０．０００５〜０．００５％の範囲内にする必要がある。好適範囲は０．００１〜０．００３％である。
【００２７】
Ａｌは脱酸剤として添加する。０．０１５％未満ではその効果は不十分である。一方、０．０５％を超えると、ＡｌＮが圧延加熱時に溶体化しないで残存し、ＮｂやＴｉ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、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０．０１５〜０．０５％の範囲内にする必要がある。好適範囲は０．０２５〜０．０４％である。
【００２８】
Ｔｉは鋼中でＮと結合してＴｉＮを生成するが、これによる固溶Ｎの固定によるＢＮの析出防止、つまり固溶Ｂの確保を目的として添加する。しかしながら、０．０３％未満ではその効果は不十分である。一方、Ｔｉを０．１５％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。好適範囲は、０．０３〜０．１３％である。
【００２９】
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。０．００２％未満ではその効果は不十分である。一方、０．０５％を超えると、素材の硬さが硬くなって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０．００２〜０．０５％の範囲内にする必要がある。好適範囲は、０．００５〜０．０３％である。
【００３０】
Ｎは以下の２点の理由から極力制限することが望ましい。▲１▼Ｂは上記のように焼入れ性向上、粒界強化等を目的として添加するが、これらのＢの効果は鋼中で固溶Ｂの状態で初めて効果を発現するため、Ｎ量を低減してＢＮの生成を抑制することが必須である。▲２▼また、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ＮｂＣ、ＮｂＣ主体のＮｂ（ＣＮ）とＴｉＣ、ＴｉＣ主体のＴｉ（ＣＮ）の析出サイトとなり、これらのＴｉの炭窒化物、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０．００５％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０．００５％未満にする必要がある。
【００３１】
次に、本発明では必要に応じて、Ｍｏ、Ｎｉの１種または２種を含有する。
【００３２】
Ｍｏは鋼に強度、焼入れ性を与えるのに有効な元素であるが、０．３％を超えて添加すると硬さの上昇を招き冷間加工性が劣化する。以上の理由から、Ｍｏは０．３％以下としたが、その含有量を０．０２〜０．３％の範囲内にすることが好ましい。
【００３３】
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、２．５％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。したがって、Ｎｉは２．５％以下としたが、その含有量を０．１〜２．５％の範囲内にすることが好ましい。
【００３４】
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（０％を含む）に制限する必要がある。好適範囲は０．０１５％以下である。
【００３５】
また、Ｏは鋼中でＡｌ２Ｏ３のような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、Ｎｂの析出物、Ｔｉの析出物の析出サイトとなり、熱間加工時にＮｂの析出物、Ｔｉ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、Ｏ量はできるだけ低減することが望ましい。図１にＯ量と結晶粒粗大化温度との関係を示す。圧下率５０％の据え込みを行った後、各温度で５時間保定して浸炭シミュレーションを行った結果である。Ｏ含有量が０．００２５％を超えると粗大粒発生温度が９５０℃以下になり、実用的には粗大粒の発生が懸念される。以上の理由から、その含有量を０．００２５％以下（０％を含む）に制限する必要がある。好適範囲は０．００２％以下である。
【００３６】
次に本発明では、熱間圧延後のマトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上を有するが、このように限定した理由を以下に述べる。
【００３７】
結晶粒の粗大化を抑制するためには、結晶粒界をピン止めする粒子を多量、微細に分散させることが有効であり、粒子の直径が小さいほど、また量が多いほどピン止め粒子の数が増加するため好ましい。本発明でいうＴｉの析出物はＴｉＣ、ＴｉＮ、Ｔｉ（ＣＮ）をさす。熱間加工後の直径０．２μｍ以下のＴｉの析出物の個数と結晶粒粗大化温度との関係を図２に示す。圧下率５０％の据え込みを行った後、各温度で５時間浸炭シミュレーションを行った結果である。図２からあきらかなように、直径０．２μｍ以下のＴｉの微細析出物をその合計で１０個／μｍ２以上分散させると実用上の浸炭加熱温度域において結晶粒の粗大化が生じず、優れた結晶粒粗大防止特性が得られる。以上から、マトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上分散していることが必要である。好適範囲は２０個／μｍ２以上である。
【００３８】
次に本発明では、熱間加工材の硬さＨＶを下記式で定義する硬さ指数ＨでＨ＋３０以下の範囲に制限するが、このように限定した理由を以下に述べる。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
【００３９】
本発明の請求項１または２では、浸炭時の粗大粒を防止するために、ＴｉＣ主体のＴｉの炭窒化物またはさらにＮｂＣ主体のＮｂの炭窒化物を浸炭時に微細分散させることを特徴としている。これらのＴｉ、Ｎｂの炭窒化物の大部分は鋼片圧延の冷却過程で析出し、その後の棒鋼線材圧延の加熱時に一部が固溶し、析出物がオストワルド成長する。析出物がオストワルド成長すると浸炭時の粗大粒防止特性は劣化する。また、棒鋼線材圧延の加熱時にＴｉ、Ｎｂが一部固溶すると、棒鋼線材圧延後の冷却過程で、オーステナイトからフェライト変態時に、Ｔｉ、Ｎｂの炭窒化物が相界面析出し、これによる析出硬化により硬さが増加する。逆に言うと、棒鋼線材圧延材の硬さは、棒鋼線材圧延の加熱時に固溶するＴｉ、Ｎｂの炭窒化物の量の程度を反映しており、棒鋼線材圧延材の硬さが硬いほど、棒鋼線材圧延の加熱時に固溶するＴｉ、Ｎｂの炭窒化物の量が多く、析出物のオストワルド成長が顕著になり、その後の浸炭時の粗大粒防止特性は劣化する。上記の理由から合金元素（Ｔｉ、Ｎｂを除く）に応じて棒鋼線材圧延材の硬さの上限値を制限することにより、棒鋼線材圧延時の冷却過程でのＴｉ、Ｎｂによる析出硬化を小さくすることができ、これにより、浸炭時のＴｉ、Ｎｂの析出物の微細分散が可能になり、浸炭時の粗大粒の防止が可能になる。さらに、鋼材の硬さの上限値を制限することにより、圧延ままでの冷間加工性は向上する。以上の技術思想から、Ｔｉ、Ｎｂを除く成分系によって決まる硬さ指数を導入し、熱間加工材の硬さの上限値を規定した。硬さ指数Ｈは、熱間加工材の硬さに及ぼす合金成分の影響を定式化した指数であり、単位はＨＶである。硬さ指数ＨにはＴｉ、Ｎｂは含まれていない。つまり、繰り返しになるが、本願発明の規定を満たす鋼材においては、棒鋼線材圧延による冷却過程でのＴｉ、Ｎｂによる析出硬化量が実質的に小さいことを意味している。また、硬さ指数Ｈを定義した前提条件として、熱間加工材にベイナイト組織が実質的に含まれないことを前提としている。
【００４０】
熱間加工材の硬さがＨＶでＨ＋３０を超えると熱間加工材の硬さが硬くなり冷間加工性が劣化するとともに浸炭時の粗大粒防止特性も劣化するので、熱間加工材の硬さをＨＶでＨ＋３０以下の範囲に制限した。好適範囲は、Ｈ−２０〜Ｈ＋２５の範囲である。
【００４１】
なお、本発明で規定する硬さ（ＨＶ）は、熱間加工材の表面脱炭層を除く最表層の硬さである。
【００４２】
次に本願発明では、粗大粒防止の目的で、脱炭深さの上限を規定している。本要件は、本発明の技術の最も重要な特徴である。表１に脱炭深さと浸炭粗大粒発生温度の関係を示す。粗大粒発生温度は、圧下率５０％の据え込みを行った後、各温度で５時間浸炭シミュレーションを行って求めた。本出願人は、脱炭深さが、ＤＭ−Ｔ０．２ｍｍを超えると、浸炭時に粗大粒が発生しやすくなることを初めて発見した。これは、浸炭加熱の昇温時に表層の脱炭部から混粒が生じ、これが粗大粒成長のきっかけになるためである。以上の理由から、脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下に制限する。このような脱炭深さは後述する低温加熱圧延を行うことによって達成できる。
【００４３】
【表１】
次に熱間圧延条件について説明する。
【００４５】
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、鋳造後、Ａ３点以下に冷却することなく分塊圧延工程を経て、線材または棒鋼に以下の低温加熱圧延を行う。
【００４６】
即ち、加熱温度は９００〜１０７０℃のＡｒ３点直上の温度として、熱間圧延の仕上げ温度を８００〜９７０℃の熱間圧延を行う。熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する条件で線材または棒鋼に熱間加工する。
【００４７】
加熱温度を９００〜１０７０℃のＡｒ３点直上の温度とするのは、鋳造後にＡ３点温度以下に冷却することなしに分塊圧延することによって生成した微細ＴｉＣ析出物をマトリックスに固溶させないようにするためであり、９００℃未満では圧延温度がフェライト域圧延となるので好ましくなく、また１０７０℃を超えると析出物がマトリックスに固溶し、微細ＴｉＣ析出物の数が減るので好ましくない。また、加熱温度が１０７０℃を超えると全脱炭が顕著になり、この点からも粗大粒防止特性は劣化する。以上のように、微細ＴｉＣ析出物を微細なままの状態で保持することにより、浸炭時に粗大粒の発生を抑制することができるようにするため、加熱温度を９００〜１０７０℃とした。
【００４８】
次に、熱間圧延の仕上げ温度を８００〜９７０℃とするのは次の理由による。仕上げ温度が８００℃未満では、圧延材のフェライト脱炭が進行するために、結果的に全脱炭も顕著になり、浸炭時に粗大粒が発生しやすくなる。一方、仕上げ温度が９７０℃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８００〜９７０℃とする。好適温度は８５０〜９６０℃である。
【００４９】
次に、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷するのは次の理由による。冷却速度が１℃／秒を超えると、ベイナイトの組織分率が大きくなり、浸炭時に粗大粒が発生しやすくなる。さらに、ベイナイトの組織分率が大きくなると、圧延材の硬さが上昇し冷間鍛造性が劣化する。そのため、冷却速度１℃／秒以下に制限する。好適範囲は０．７℃／秒以下である。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。
【００５０】
本発明では、鋳片のサイズ、凝固時の冷却速度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。また、本発明鋼は、圧延ままの棒鋼を冷間鍛造で部品に成形する工程だけでなく、冷間鍛造の前に焼鈍工程や温・熱間鍛造を経由する場合、温・熱間鍛造工程で部品に成形される場合、切削工程で部品に成形される場合にも適用できる。
【００５１】
【実施例】
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。
【００５２】
表２に示す組成を有する転炉溶製鋼を連続鋳造し、鋳造後、鋼をＡ３点温度以下に冷却することなく分塊圧延を行い、１６２ｍｍ角の鋼片（圧延素材）とした（分塊圧延条件Ｉ）。比較鋼ａ、ｂについては、連続鋳造後、鋼を一旦常温まで冷却した後、再度Ａ３点以上に加熱して分塊圧延を行い、１６２ｍｍ角の鋼片（圧延素材）とした（分塊圧延条件ＩＩ）。
【００５３】
続いて、熱間加工により、直径３４ｍｍの棒鋼を製造した。熱間加工条件を表３に示す。熱間加工後の冷却速度は冷却床に設置した徐冷カバーを用いて調整した。比較鋼Ｙ、ＺはＪＩＳのＳＣｒ４２０およびＳＣＭ４２０である。
【００５４】
熱間加工後の棒鋼のＴｉの析出物、Ｎｂの析出物の分散状態を調べるために、棒鋼のマトリックス中に存在する析出物を抽出レプリカ法によって採取し、透過型電子顕微鏡で観察した。観察方法は３００００倍で２０視野程度観察し、１視野中の直径０．２μｍ以下のＴｉの析出物、Ｎｂの析出物、ＴｉとＮｂの複合組成からなる析出物の数を数え、１平方μｍ当たりの数に換算した。
【００５５】
圧延後の棒鋼のビッカース硬さを測定した。また、ミクロ観察、全脱炭深さの調査も行った。さらに、圧延ままの棒鋼から、据え込み試験片を作成し、冷間加工性の指標として、冷間変形抵抗と限界圧縮率を求めた。冷間変形抵抗は相当歪み１．０における変形抵抗で代表させた。
【００５６】
次に、圧延ままの棒鋼から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大粒発生温度（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、粗大粒発生温度が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大粒発生と判定した。
【００５７】
さらに、直径３０ｍｍの棒鋼を削り出し、直径２２ｍｍへ引き抜きを行った後、９４０℃×４時間の条件で浸炭焼き入れを行い、γ粒度を測定した。
【００５８】
これらの調査結果を熱間加工条件とあわせて表３に示す。
【００５９】
比較例２５、２６はＪＩＳのＳＣｒ４２０およびＳＣＭ４２０の特性であるが、本発明例の冷間変形抵抗は、比較例２５、２６に比較して顕著に小さく、また限界据え込み率も優れている。また、本発明例の結晶粒粗大化温度は９７０℃以上であり、通常の上限の浸炭条件である９５０℃では、粗大粒の発生を防止できることが明らかである。
【００６０】
次に、表３において、比較例１９はＳｉの含有量が本願規定の範囲を上回った場合であり、本発明例に比較して、冷間加工性は劣る。
【００６１】
比較例２０はＴｉの含有量が本願規定の範囲を下回った場合であり、比較例２１はＮの含有量が本願規定の範囲を上回った場合であり、比較例２２はＯの含有量が本願規定の範囲を上回った場合であり、いずれも粗大粒防止特性は劣っている。
【００６２】
比較例２３はＴｉの含有量が本願規定の範囲を上回った場合であり、比較例２４はＮｂの含有量が本願規定の範囲を上回った場合であり、いずれも本発明例に比較して、冷間加工性は劣る。
【００６３】
比較例２７，２８は鋼片の製造方法が本願発明と異なり、鋳造後、鋼をＡ３点温度以下に一旦冷却した後分塊圧延を行う方法で製造した場合であり、いずれも粗大粒防止特性は劣っている。
【００６４】
比較例２９は熱間加工の加熱温度が本願規定の範囲を上回った場合であり、析出物個数は本願発明の範囲を下回り、圧延後の硬さは本願規定の範囲を上回り、全脱炭深さも本願発明の範囲を上回り、粗大粒防止特性は劣っている。
【００６５】
比較例３０は熱間加工の仕上げ温度が本願規定の範囲を上回った場合であり、本発明例４に比較して冷間加工性は劣る。比較例３１は熱間加工の仕上げ温度が本願規定の範囲を下回った場合であり、全脱炭深さは本願発明の範囲を上回り、粗大粒防止特性は劣っている。比較例３２は熱間加工後の冷却速度が本願規定の範囲を上回った場合であり、粗大粒防止特性が劣るとともに、冷間加工性も劣る。
【００６６】
【表２】
【表３】
【発明の効果】
本発明の冷間加工性と低浸炭歪み特性に優れた肌焼鋼とその製造方法を用いれば、冷間鍛造時には冷間加工性に優れ、同時に冷間鍛造工程で製造しても、浸炭時に粗大粒の発生を安定的に抑制することができ、これにより、歪みや曲がりの発生を防止することができる。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り、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。
【図面の簡単な説明】
【図１】Ｏ量と結晶粒粗大化温度との関係を示す図である。
【図２】熱間加工後の直径０．２μｍ以下のＴｉの析出物の個数と結晶粒粗大化温度との関係を示す図である。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a case-hardening steel material excellent in cold workability and excellent in preventing coarse grains during carburizing, and a method for producing the same.
[0002]
[Prior art]
Gears, shafts, and CVJ parts are usually made of medium carbon alloy steel for machine structural use as defined in JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc., and cold forging (also rolling) -It is manufactured in a process of carburizing and quenching after being processed into a predetermined shape by 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. Here, when switching from hot forging to cold forging, it is important to reduce the cold deformation resistance of steel and to improve the critical compression ratio. This is because the former is for ensuring the life of the forging tool, and the latter is for preventing cracking of the steel during cold forging. As a steel material suitable for such cold forging, application of boron steel has been studied. However, boron steel tends to cause a phenomenon in which some austenite crystal grains become coarse during carburizing heating. Parts with such coarse particles generate heat treatment distortion after carburizing and quenching. For example, gears and CVJ parts with large carburizing distortion may cause noise and vibration. Under these circumstances, there is a strong demand for boron steel that has excellent cold workability and excellent low carburizing strain characteristics, that is, does not generate coarse grains during carburizing.
[0003]
The present applicant has previously provided an invention of a case-hardened boron steel excellent in cold workability and capable of preventing the generation of coarse grains during carburization and excellent in low carburizing distortion characteristics and a method for producing the same (Patent Document 1). .
[0004]
In this invention, coarse grains cannot be prevented only by limiting the chemical composition of the steel. Therefore, the solid solution state of AlN in the steel material after hot working and fine precipitates such as Nb (CN) and Ti (CN) A means of controlling the amount of precipitation and utilizing NbC and TiC precipitates as pinning particles is adopted.
[0005]
That is, the rolling heating temperature is set to a high temperature of 1150 ° C. or higher, the precipitation amount of AlN is limited in the state of the steel material after hot working, and the Nb precipitate is set to a high heating temperature of 1150 ° C. or higher. The Ti precipitates are once in solution, and after the hot rolling, the precipitation temperature range of these precipitates is gradually cooled, so that a certain amount of Nb precipitates and Ti precipitates are added to the steel after hot rolling. A large amount is finely dispersed.
[0006]
In addition, the present applicant is excellent in cold workability during cold forging, and can prevent generation of coarse grains during carburization and formation of an incompletely quenched structure formed to a depth of 0.2 to 0.7 mm from the surface. The invention of the case hardening boron steel which can be performed, and its manufacturing method was provided (patent document 2).
[0007]
In this invention, the heating temperature of steel with a limited chemical composition is set to a high temperature of 1150 ° C. or higher, the finishing temperature of hot rolling is 840 to 1000 ° C., and the temperature range of 800 to 500 ° C. is 1 ° C. / It is characterized by slow cooling at a cooling rate of less than a second.
[0008]
In any of the above inventions, the heating temperature for rolling is set to a high temperature of 1150 ° C. or higher, and Ti precipitates and Nb precipitates are once solutionized. Finely dispersed. However, even in these methods, the prevention of coarse grains during carburizing may not always be sufficient, and when the rolling heating temperature is set to a high temperature of 1150 ° C. or higher, the rolling becomes hard and the cold workability becomes insufficient. In addition, there is a problem that total decarburization becomes prominent.
[0009]
[Patent Document 1]
Japanese Patent Laid-Open No. 11-335777 [Patent Document 2]
JP 2001-303172 A [Problems to be solved by the invention]
It is an object of the present invention to provide a case-hardening steel material capable of improving cold workability and preventing the generation of coarse grains during carburizing and a method for producing the same.
[0010]
[Means to solve the problem]
The present inventor has intensively studied to improve the cold workability of the case-hardening boron steel and to improve the characteristics of preventing coarse grains during carburizing. As a result, when the rolling heating temperature is set to a high temperature of 1150 ° C. or higher as in the conventional case, when the total decarburization becomes significant, coarse particles are generated from the decarburized portion at the time of subsequent carburization, and thus the coarse particle prevention characteristics are insufficient. I found out. Further, the rolling heating temperature is set to a high temperature of 1150 ° C. or higher as in the prior art, and even if the Ti precipitate and the Nb precipitate are not once solutionized, the HCR is divided without cooling to the A3 point temperature or less after casting. By rolling the rolled steel slab at low temperature, it is possible to reduce total decarburization, thereby preventing the formation of coarse grains and improving cold workability by softening by reducing the precipitation strengthening amount. The present invention has been completed by finding out what can be done.
[0011]
The gist of the present invention is as follows.
[0012]
(1) In mass%,
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
Containing the balance being iron and unavoidable impurities, hot precipitates diameter 0.2μm following Ti in the matrix of the tissue after rolling has 10 / [mu] m 2 or more, the hardness index H When defined by the following formula, the hardness is HV and H + 30 or less, the decarburization depth specified by JISG0558: DM-T 0.2 mm or less, cold workability and coarse grain prevention characteristics during carburizing Excellent steel for case hardening.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0013]
(2) By mass%
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
Nb: 0.002 to 0.05%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
In which the balance consists of iron and inevitable impurities, and the precipitate is composed of Nb precipitates having a diameter of 0.2 μm or less, Ti precipitates, or a composite composition of Nb and Ti in the matrix of the structure after hot rolling. The total of the products has 10 pieces / μm 2 or more, and when the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, and the decarburization depth specified by JISG0558: DM-T 0.2 mm A case-hardening steel excellent in cold workability and preventing coarse grains during carburizing, characterized by the following:
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0014]
(3) Furthermore, in mass%,
Mo: 0.3% or less,
Ni: Skin hardening excellent in cold workability and coarse grain prevention property at the time of carburizing described in (1) or (2) above, containing one or two of 2.5% or less Steel material.
[0015]
(4) By mass%
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
Steel, the balance of which is made of iron and inevitable impurities, and a steel slab manufactured by a step of performing a batch rolling without cooling to below the A3 point temperature after casting, a heating temperature of 900 to 1070 ° C., heat Hot rolling to 800 to 970 ° C. and hot rolling to wire or bar steel under conditions where the temperature range from 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less. When the number of Ti precipitates having a diameter of 0.2 μm or less is 10 / μm 2 or more in the matrix of the structure after rolling and the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, according to JISG0558 Decarburization depth to be defined: DM-T 0.2 mm or less A method for producing a case-hardening steel material excellent in cold workability and preventing coarse grains during carburizing.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0016]
(5) By mass%
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
Nb: 0.002 to 0.05%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
Steel, the balance of which is made of iron and inevitable impurities, and a steel slab manufactured by a step of performing a batch rolling without cooling to below the A3 point temperature after casting, a heating temperature of 900 to 1070 ° C., heat Hot rolling to 800 to 970 ° C. and hot rolling to wire or bar steel under conditions where the temperature range from 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less. In the matrix of the structure after rolling, a total of 10 N / μm 2 of Nb precipitates having a diameter of 0.2 μm or less, Ti precipitates, or precipitates composed of a composite composition of Nb and Ti, When the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, and the decarburization depth specified by JISG0558: DM-T 0.2 mm or less. Excellent coarse grain prevention properties Method of manufacturing a baked for steel.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0017]
(6) Furthermore, in mass%,
Mo: 0.3% or less,
Ni: Skin hardening excellent in cold workability and coarse grain prevention property at the time of carburizing described in (4) or (5) above, containing one or two of 2.5% or less Steel manufacturing method.
[0018]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be described in detail.
[0019]
In the present invention, when steel strip produced by HCR (Hot Charge Rolling) is used to produce steel pieces that have been subjected to lump rolling without being cooled to A3 point temperature or lower after steel casting, when steel bar wire rolling is performed at low temperature and hot rolling, fine precipitation of TiC A thing is kept in the state which precipitated as it is. The precipitation strengthening amount is lower than the precipitation strengthening amount precipitated from the solid solution at a high temperature. Further, the total decarburization amount is also reduced, and the decarburization depth specified in JIS G 0558: DM-T can be 0.2 mm or less. As a result, the generation of coarse grains can be prevented, and the cold workability can be improved by softening.
[0020]
First, the reasons for limiting the components will be described.
[0021]
C is an effective element for imparting the necessary strength to the steel, but if it is less than 0.1%, the necessary tensile strength cannot be secured, and if it exceeds 0.3%, it becomes hard and cold work is performed. The core portion toughness after carburization is deteriorated as well as the property is deteriorated, so it is necessary to be within the range of 0.1 to 0.3%.
[0022]
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 0.01%, the effect is ineffective. It is enough. On the other hand, if it exceeds 0.15%, the hardness is increased and the cold forgeability is deteriorated. For the above reasons, the content needs to be in the range of 0.01 to 0.15%. The preferred range is 0.02 to 0.1%.
[0023]
Mn is an element effective for deoxidation of steel and is an element effective for imparting necessary strength and hardenability to the steel, but if less than 0.2%, the effect is insufficient, 0.65% If it exceeds, the effect is not only saturated, but also the hardness is increased and the cold forgeability deteriorates, so it is necessary to be within the range of 0.2 to 0.65%. The preferred range is 0.3-0.6%.
[0024]
S forms MnS in the steel and is added for the purpose of improving machinability. However, if it is less than 0.005%, 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 0.005 to 0.15%. The preferred range is 0.005 to 0.04%.
[0025]
Cr is an element effective for imparting strength and hardenability to steel, but if it is less than 0.4%, the effect is insufficient, and if added over 1.25%, it causes an increase in hardness and cold. Forgeability deteriorates. For the above reasons, the content needs to be in the range of 0.4 to 1.25%. The preferred range is 0.6-1.0%.
[0026]
B is added for the following three points. (1) In steel bar / wire rolling, boron iron carbide is generated in the cooling process after rolling, thereby increasing the growth rate of ferrite and promoting softening during rolling. (2) When carburizing and quenching, impart hardenability to the steel. (3) Improve the fatigue strength and impact strength of carburized parts by improving the grain boundary strength of the carburized material. When the amount is less than 0.0005%, the above effect is insufficient. When the amount exceeds 0.005%, the effect is saturated, so the content needs to be within the range of 0.0005 to 0.005%. There is. The preferred range is 0.001 to 0.003%.
[0027]
Al is added as a deoxidizer. If it is less than 0.015%, the effect is insufficient. On the other hand, if it exceeds 0.05%, AlN remains without being melted during rolling and heating, and becomes a precipitation site of Nb and Ti precipitates, which inhibits fine dispersion of these precipitates and coarsens the crystal grains. To encourage. For the above reasons, the content needs to be in the range of 0.015 to 0.05%. The preferred range is 0.025 to 0.04%.
[0028]
Ti combines with N in the steel to produce TiN, and is added for the purpose of preventing precipitation of BN by fixing solid solution N, that is, securing solid solution B. However, if it is less than 0.03%, the effect is insufficient. On the other hand, if Ti is added in an amount exceeding 0.15%, precipitation hardening due to TiC becomes remarkable, and cold workability is remarkably deteriorated. The preferred range is 0.03 to 0.13%.
[0029]
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. If it is less than 0.002%, the effect is insufficient. On the other hand, if it exceeds 0.05%, the hardness of the material becomes hard and the cold forgeability deteriorates, and it becomes difficult to form a solution during heating of the steel bar and wire rod. For the above reasons, the content needs to be in the range of 0.002 to 0.05%. The preferred range is 0.005 to 0.03%.
[0030]
It is desirable to limit N as much as possible for the following two reasons. (1) B is added for the purpose of improving hardenability and strengthening grain boundaries as described above. However, since the effect of these B is manifested only in the state of solid solution B in steel, the amount of N is reduced. Therefore, it is essential to suppress the generation of BN. (2) When N is combined with Ti in steel, it produces coarse TiN that hardly contributes to grain control. This is the precipitation site of NbC, NbC-based Nb (CN) and TiC, TiC-based Ti (CN). Thus, the fine precipitation of these Ti carbonitrides and Nb carbonitrides is inhibited, and the formation of coarse particles is promoted. The above adverse effect is particularly remarkable when the N amount is 0.005% or more. For the above reasons, the content needs to be less than 0.005%.
[0031]
Next, in this invention, 1 type or 2 types of Mo and Ni are contained as needed.
[0032]
Mo is an element effective for imparting strength and hardenability to steel, but if added over 0.3%, the hardness increases and cold workability deteriorates. For these reasons, the Mo content is set to 0.3% or less, but the content is preferably within the range of 0.02 to 0.3%.
[0033]
Ni is also an effective element for imparting strength and hardenability to steel, but if added over 2.5%, the hardness is increased and cold forgeability is deteriorated. Therefore, although Ni was made into 2.5% or less, it is preferable to make the content into the range of 0.1-2.5%.
[0034]
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 (including 0%). The preferred range is 0.015% or less.
[0035]
O forms oxide inclusions such as Al 2 O 3 in the steel. When a large amount of oxide inclusions are present in the steel, it becomes a precipitation site for Nb precipitates and Ti precipitates, Nb precipitates and Ti precipitates are coarsely deposited during hot working, and crystallized during carburization. Grain coarsening cannot be suppressed. Therefore, it is desirable to reduce the amount of O as much as possible. FIG. 1 shows the relationship between the amount of O and the crystal grain coarsening temperature. This is the result of carburizing simulation after holding up at a rolling reduction of 50% and holding at each temperature for 5 hours. When the O content exceeds 0.0025%, the coarse particle generation temperature becomes 950 ° C. or less, and there is a concern that the generation of coarse particles is practically concerned. For the above reasons, the content needs to be limited to 0.0025% or less (including 0%). The preferred range is 0.002% or less.
[0036]
Next, in the present invention, the hot-rolled matrix has 10 precipitates / μm 2 or more of Ti having a diameter of 0.2 μm or less. The reason for this limitation will be described below.
[0037]
In order to suppress the coarsening of crystal grains, it is effective to disperse a large amount and finely the particles that pin the crystal grain boundaries. The smaller the diameter of the particles and the larger the amount, the more pinned particles. Is preferable because of an increase. The Ti precipitate referred to in the present invention refers to TiC, TiN, and Ti (CN). FIG. 2 shows the relationship between the number of Ti precipitates having a diameter of 0.2 μm or less after hot working and the crystal grain coarsening temperature. This is a result of performing a carburization simulation for 5 hours at each temperature after upsetting at a rolling reduction of 50%. As is clear from FIG. 2, when fine precipitates of Ti having a diameter of 0.2 μm or less are dispersed in a total of 10 particles / μm 2 or more, the coarsening of crystal grains does not occur in a practical carburizing heating temperature range, which is excellent. The crystal grain coarseness preventing property is obtained. From the above, it is necessary to disperse 10 precipitates / μm 2 or more of Ti having a diameter of 0.2 μm or less in the matrix. A preferable range is 20 pieces / μm 2 or more.
[0038]
Next, in the present invention, the hardness HV of the hot-worked material is limited to a range of H + 30 or less with a hardness index H defined by the following formula. The reason for this limitation will be described below.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0039]
Claim 1 or 2 of the present invention is characterized in that TiC-based Ti carbonitride or further NbC-based Nb carbonitride is finely dispersed during carburizing in order to prevent coarse grains during carburizing. . Most of these Ti and Nb carbonitrides precipitate in the cooling process of the steel slab rolling, and partly dissolve in the subsequent heating of the bar steel wire rolling, and the precipitates grow Ostwald. When the precipitate grows through Ostwald, the coarse grain prevention characteristics during carburization deteriorate. Also, if Ti and Nb are partly dissolved during heating of the steel bar wire rolling, Ti and Nb carbonitrides precipitate at the interfacial phase during the transformation of austenite to ferrite in the cooling process after the steel bar wire rolling, and precipitation hardening occurs thereby. Increases the hardness. In other words, the hardness of the rolled steel bar reflects the amount of Ti and Nb carbonitrides that are dissolved during heating of the rolled steel bar, and the harder the rolled steel bar is, the harder it is. In addition, the amount of Ti and Nb carbonitrides that are solid-dissolved during heating of the steel rod rolling is large, and the Ostwald growth of precipitates becomes significant, and the coarse grain prevention characteristics during subsequent carburizing deteriorate. For the above reasons, by limiting the upper limit value of the hardness of the rolled steel bar according to the alloying elements (excluding Ti and Nb), the precipitation hardening due to Ti and Nb in the cooling process at the time of rolling the steel bar is reduced. This makes it possible to finely disperse precipitates of Ti and Nb at the time of carburizing, and to prevent coarse particles at the time of carburizing. Furthermore, by limiting the upper limit value of the hardness of the steel material, the cold workability as it is rolled is improved. From the above technical idea, the hardness index determined by the component system excluding Ti and Nb was introduced to define the upper limit value of the hardness of the hot-worked material. The hardness index H is an index that formulates the influence of the alloy component on the hardness of the hot-worked material, and its unit is HV. The hardness index H does not contain Ti or Nb. That is, although it repeats, in the steel material which satisfy | fills the prescription | regulation of this invention, it means that the precipitation hardening amount by Ti and Nb in the cooling process by steel bar wire rod rolling is substantially small. As a precondition for defining the hardness index H, it is assumed that the hot-worked material does not substantially contain a bainite structure.
[0040]
If the hardness of the hot work material is HV and exceeds H + 30, the hardness of the hot work material becomes hard and the cold workability deteriorates, and the coarse grain prevention property at the time of carburizing also deteriorates. The thickness was limited to a range of H + 30 or less in HV. The preferred range is H-20 to H + 25.
[0041]
In addition, the hardness (HV) prescribed | regulated by this invention is the hardness of the outermost layer except the surface decarburization layer of a hot work material.
[0042]
Next, in this invention, the upper limit of the decarburization depth is prescribed | regulated for the purpose of coarse grain prevention. This requirement is the most important feature of the technology of the present invention. Table 1 shows the relationship between the decarburization depth and the carburized coarse particle generation temperature. The coarse grain generation temperature was determined by performing a carburization simulation for 5 hours at each temperature after upsetting at a rolling reduction of 50%. The present applicant has discovered for the first time that when the decarburization depth exceeds DM-T 0.2 mm, coarse grains are likely to be generated during carburizing. This is because mixed grains are generated from the decarburized portion of the surface layer at the time of raising the temperature of the carburizing heating, and this is a trigger for coarse grain growth. For the above reasons, the decarburization depth is limited to DM-T 0.2 mm or less. Such a decarburization depth can be achieved by performing low-temperature heating rolling described later.
[0043]
[Table 1]
Next, hot rolling conditions will be described.
[0045]
The steel of the present invention is melted by a usual method such as a converter, electric furnace, etc., the components are adjusted, and after casting, it is subjected to a block rolling process without cooling to the A3 point or lower, to a wire rod or bar steel. The following low temperature heating rolling is performed.
[0046]
That is, hot rolling is performed at a finishing temperature of 800 to 970 ° C. with a heating temperature of 900 to 1070 ° C. just above the Ar 3 point. Subsequent to hot rolling, the wire or steel bar is hot worked under the condition that the temperature range of 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less.
[0047]
To the heating temperature and the temperature just above the Ar 3 point of 900 to 1,070 ° C. is not to fine TiC precipitates produced by slabbing without cooling after the A3 point temperature below the casting form a solid solution in the matrix If the temperature is lower than 900 ° C., the rolling temperature becomes ferrite region rolling, which is not preferable, and if it exceeds 1070 ° C., the precipitates are dissolved in the matrix and the number of fine TiC precipitates is decreased. Further, when the heating temperature exceeds 1070 ° C., total decarburization becomes remarkable, and from this point, the coarse grain prevention characteristics deteriorate. As described above, the heating temperature is set to 900 to 1070 ° C. in order to suppress the generation of coarse particles during carburization by holding the fine TiC precipitate in a fine state.
[0048]
Next, the finishing temperature of hot rolling is set to 800 to 970 ° C. for the following reason. If the finishing temperature is less than 800 ° C., ferrite decarburization of the rolled material proceeds, and as a result, total decarburization becomes prominent, and coarse particles are likely to be generated during carburizing. On the other hand, when the finishing temperature exceeds 970 ° C., the hardness of the rolled material becomes hard and the cold forgeability deteriorates. For the above reasons, the hot rolling finishing temperature is set to 800 to 970 ° C. The preferred temperature is 850 to 960 ° C.
[0049]
Next, following the hot rolling, the temperature range of 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less for the following reason. When the cooling rate exceeds 1 ° C./second, the structure fraction of bainite increases, and coarse grains are likely to be generated during carburizing. Furthermore, when the structure fraction of bainite increases, the hardness of the rolled material increases and the cold forgeability deteriorates. Therefore, the cooling rate is limited to 1 ° C./second or less. The preferred range is 0.7 ° C./second 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.
[0050]
In the present invention, the size of the slab and the cooling rate during solidification are not particularly limited, and any condition may be used as long as the requirements of the present invention are satisfied. In addition, the steel of the present invention is not only a process for forming an as-rolled steel bar into parts by cold forging, but also a warm / hot forging process when passing through an annealing process or warm / hot forging before cold forging. It can also be applied to the case where it is formed into a part by the cutting process.
[0051]
【Example】
Hereinafter, the effects of the present invention will be described more specifically by way of examples.
[0052]
Converter molten steel having the composition shown in Table 2 is continuously cast, and after casting, the steel is subjected to partial rolling without cooling to a temperature below the A3 point temperature to obtain a 162 mm square steel slab (rolling material). Rolling conditions I). For the comparative steels a and b, after continuous casting, the steel was once cooled to room temperature, and then heated again to the A3 point or higher and subjected to split rolling to obtain a 162 mm square steel slab (rolling material). Condition II).
[0053]
Subsequently, a steel bar having a diameter of 34 mm was manufactured by hot working. Table 3 shows the hot working conditions. The cooling rate after hot working was adjusted using a slow cooling cover installed on the cooling bed. The comparative steels Y and Z are JIS SCr420 and SCM420.
[0054]
In order to investigate the dispersion state of Ti precipitates and Nb precipitates in the steel bar after hot working, the precipitates present in the bar steel matrix were collected by the extraction replica method and observed with a transmission electron microscope. The observation method is to observe about 20 fields of view at 30000 times, and count the number of Ti precipitates having a diameter of 0.2 μm or less in one field of view, Nb precipitates, and precipitates composed of a composite composition of Ti and Nb per square μm. Converted to a number.
[0055]
The Vickers hardness of the rolled steel bar was measured. Microscopic observation and total decarburization depth were also investigated. Furthermore, an upsetting test piece was prepared from the rolled steel bar, and the cold deformation resistance and the critical compressibility were obtained as indicators of cold workability. The cold deformation resistance was represented by the deformation resistance at an equivalent strain of 1.0.
[0056]
Next, an upsetting test piece was prepared from the rolled steel bar, and after upsetting with a reduction ratio of 50%, carburization 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., those having a coarse grain generation temperature of 950 ° C. or less were determined to be inferior in crystal grain coarsening characteristics. The prior austenite grain size was measured in accordance with JIS G 0551, observed at 400 magnifications for about 10 fields of view, and if any coarse grain having a grain size number of 5 or less was present, it was determined that coarse grains were generated.
[0057]
Furthermore, after cutting out a steel bar having a diameter of 30 mm and drawing it to a diameter of 22 mm, carburizing and quenching was performed under conditions of 940 ° C. × 4 hours, and the γ particle size was measured.
[0058]
These survey results are shown in Table 3 together with hot working conditions.
[0059]
Comparative examples 25 and 26 are the characteristics of JIS SCr420 and SCM420, but the cold deformation resistance of the examples of the present invention is remarkably smaller than that of comparative examples 25 and 26, and the limit upsetting rate is also excellent. Further, the crystal grain coarsening temperature of the present invention example is 970 ° C. or higher, and it is clear that the generation of coarse grains can be prevented at 950 ° C., which is a normal upper limit carburizing condition.
[0060]
Next, in Table 3, Comparative Example 19 is a case where the Si content exceeds the range specified in the present application, and the cold workability is inferior compared to the inventive example.
[0061]
Comparative Example 20 is a case where the Ti content falls below the range stipulated in the present application, Comparative Example 21 is a case where the N content exceeds the range stipulated by the present application, and Comparative Example 22 is a case where the content of O is the present level. It is a case where it exceeds the specified range, and in all cases, the coarse grain preventing property is inferior.
[0062]
Comparative Example 23 is a case where the Ti content exceeds the range specified in the present application, and Comparative Example 24 is a case where the Nb content exceeds the range specified in the present application. Cold workability is poor.
[0063]
Comparative Examples 27 and 28 differ from the present invention in the method of manufacturing the steel slab, and are the cases where the steel is manufactured by a method of performing batch rolling after the steel is once cooled to the A3 point temperature or less after casting, both of which have coarse grain prevention characteristics. Is inferior.
[0064]
Comparative Example 29 is a case where the heating temperature of hot working exceeds the range specified in the present application, the number of precipitates is below the range of the present invention, the hardness after rolling exceeds the range specified in the present application, and the total decarburization depth. Moreover, it exceeds the range of the present invention, and the coarse grain preventing property is inferior.
[0065]
Comparative Example 30 is a case where the hot working finishing temperature exceeds the range specified in the present application, and the cold workability is inferior to that of Invention Example 4. The comparative example 31 is a case where the finishing temperature of hot working falls below the range specified in the present application, the total decarburization depth exceeds the range of the present invention, and the coarse grain prevention property is inferior. The comparative example 32 is a case where the cooling rate after hot working exceeds the range specified in the present application, and the coarse grain prevention property is inferior and the cold workability is also inferior.
[0066]
[Table 2]
[Table 3]
【The invention's effect】
By using the case-hardened steel excellent in cold workability and low carburizing distortion characteristics and its manufacturing method of the present invention, it is excellent in cold workability at the time of cold forging. Generation | occurrence | production of a coarse grain can be suppressed stably and, thereby, generation | occurrence | production of a distortion and bending can be prevented. Therefore, it is possible to cold forge parts that have been difficult to cold forge due to the problem of coarse particles, and it is also possible to omit the annealing after cold forging, the industrial effect of the present invention Is extremely prominent.
[Brief description of the drawings]
FIG. 1 is a graph showing the relationship between the amount of O and the crystal grain coarsening temperature.
FIG. 2 is a diagram showing the relationship between the number of Ti precipitates having a diameter of 0.2 μm or less after hot working and the crystal grain coarsening temperature.
Claims (6)
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなり、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％% By mass
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
The balance consists of iron and inevitable impurities, and has a precipitate of Ti having a diameter of 0.2 μm or less in a matrix of a structure after hot rolling of 10 pieces / μm 2 or more, and a hardness index H When defined by the following formula, the hardness is HV and H + 30 or less, the decarburization depth specified by JISG0558: DM-T 0.2 mm or less, cold workability and coarse grain prevention characteristics during carburizing Excellent steel for case hardening.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎｂ：０．００２〜０．０５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなり、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＮｂの析出物、Ｔｉの析出物、またはＮｂとＴｉの複合組成からなる析出物のその合計で１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％% By mass
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
Nb: 0.002 to 0.05%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
In which the balance consists of iron and inevitable impurities, and the precipitate is composed of Nb precipitates having a diameter of 0.2 μm or less, Ti precipitates, or a composite composition of Nb and Ti in the matrix of the structure after hot rolling. The total of the products has 10 pieces / μm 2 or more, and when the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, and the decarburization depth specified by JISG0558: DM-T 0.2 mm A case-hardening steel excellent in cold workability and preventing coarse grains during carburizing, characterized by the following:
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
Ｍｏ：０．３％以下、
Ｎｉ：２．５％以下
のうちの１種または２種を含有することを特徴とする請求項１または２記載の冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材。Furthermore, in mass%,
Mo: 0.3% or less,
The steel for skin hardening excellent in cold workability and coarse grain prevention characteristics during carburization according to claim 1 or 2, characterized by containing one or two of Ni: 2.5% or less.
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなる鋼を鋳造後Ａ３点温度以下に冷却することなく分塊圧延を行う工程により製造された鋼片を用い、加熱温度を９００〜１０７０℃、熱間圧延の仕上温度を８００〜９７０℃、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する条件により線材または棒鋼に熱間圧延し、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＴｉの析出物を１０個／μｍ２以上とし、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％% By mass
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
Steel, the balance of which is made of iron and inevitable impurities, and a steel slab manufactured by a step of performing a batch rolling without cooling to below the A3 point temperature after casting, a heating temperature of 900 to 1070 ° C., heat Hot rolling to 800 to 970 ° C. and hot rolling to wire or bar steel under conditions where the temperature range from 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less. When the number of Ti precipitates having a diameter of 0.2 μm or less is 10 / μm 2 or more in the matrix of the structure after rolling and the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, according to JISG0558 Decarburization depth to be defined: DM-T 0.2 mm or less A method for producing a case-hardening steel material excellent in cold workability and preventing coarse grains during carburizing.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
Ｃ：０．１〜０．３％、
Ｓｉ：０．０１〜０．１５％、
Ｍｎ：０．２〜０．６５％、
Ｓ：０．００５〜０．１５％、
Ｃｒ：０．４〜１．２５％、
Ｂ：０．０００５〜０．００５％、
Ａｌ：０．０１５〜０．０５％、
Ｔｉ：０．０３〜０．１５％、
Ｎｂ：０．００２〜０．０５％、
Ｎ：０．００５％未満（０％を含む）、
Ｐ：０．０２５％以下（０％を含む）、
Ｏ：０．００２５％以下（０％を含む）
を含有し、残部が鉄および不可避的不純物からなる鋼を鋳造後Ａ３点温度以下に冷却することなく分塊圧延を行う工程により製造された鋼片を用い、加熱温度を９００〜１０７０℃、熱間圧延の仕上温度を８００〜９７０℃、熱間圧延に引き続いて８００〜５００℃の温度範囲を１℃／秒以下の冷却速度で徐冷する条件により線材または棒鋼に熱間圧延し、熱間圧延後の組織のマトリックス中に直径０．２μｍ以下のＮｂの析出物、Ｔｉの析出物、またはＮｂとＴｉの複合組成からなる析出物のその合計で１０個／μｍ２以上を有し、硬さ指数Ｈを下記式で定義すると、硬さがＨＶでＨ＋３０以下であり、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下であることを特徴とする冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％% By mass
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.65%,
S: 0.005 to 0.15%,
Cr: 0.4 to 1.25%,
B: 0.0005 to 0.005%,
Al: 0.015 to 0.05%,
Ti: 0.03-0.15%,
Nb: 0.002 to 0.05%,
N: less than 0.005% (including 0%),
P: 0.025% or less (including 0%),
O: 0.0025% or less (including 0%)
Steel, the balance of which is made of iron and inevitable impurities, and a steel slab manufactured by a step of performing a batch rolling without cooling to below the A3 point temperature after casting, a heating temperature of 900 to 1070 ° C., heat Hot rolling to 800 to 970 ° C. and hot rolling to wire or bar steel under conditions where the temperature range from 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less. In the matrix of the structure after rolling, a total of 10 N / μm 2 of Nb precipitates having a diameter of 0.2 μm or less, Ti precipitates, or precipitates composed of a composite composition of Nb and Ti, When the hardness index H is defined by the following formula, the hardness is HV and H + 30 or less, and the decarburization depth specified by JISG0558: DM-T 0.2 mm or less. Excellent coarse grain prevention properties Method of manufacturing a baked for steel.
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
Ｍｏ：０．３％以下、
Ｎｉ：２．５％以下
のうちの１種または２種を含有することを特徴とする請求項４または５記載の冷間加工性と浸炭時の粗大粒防止特性に優れた肌焼用鋼材の製造方法。Furthermore, in mass%,
Mo: 0.3% or less,
Ni: One or two of 2.5% or less is contained, The steel for skin hardening excellent in cold workability and coarse grain prevention characteristics at the time of carburizing according to claim 4 or 5 Production method.
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Cited By (12)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2008001940A (en) * | 2006-06-21 | 2008-01-10 | Kobe Steel Ltd | Method for designing component in alternative steel for chromium-molybdenum steel |
| JP2008081841A (en) * | 2006-08-28 | 2008-04-10 | Kobe Steel Ltd | Case hardening steel having excellent cold forgeability and crystal grain coarsening prevention property, and machine part obtained therefrom |
| JP4528363B1 (en) * | 2009-04-06 | 2010-08-18 | 新日本製鐵株式会社 | Case-hardened steel with excellent cold workability, machinability, and fatigue characteristics after carburizing and quenching, and method for producing the same |
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2002
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