JP3738004B2

JP3738004B2 - Case-hardening steel with excellent cold workability and prevention of coarse grains during carburizing, and its manufacturing method

Info

Publication number: JP3738004B2
Application number: JP2002372541A
Authority: JP
Inventors: 達朗越智; 学久保田; 三喜男倉地
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 2002-12-24
Filing date: 2002-12-24
Publication date: 2006-01-25
Anticipated expiration: 2022-12-24
Also published as: JP2004204263A

Description

【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及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ＪＩＳＧ４０５２、ＪＩＳＧ４１０４、ＪＩＳＧ４１０５、ＪＩＳＧ４１０６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。
【０００３】
肌焼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。
【０００４】
従来は、浸炭処理におけるオーステナイト結晶粒の粗大化を防止するために、熱間圧延後の鋼材中のＡｌＮの存在形態や、Ｎｂ（ＣＮ）、Ｔｉ（ＣＮ）等の微細析出物の析出量等を制御して、ＮｂＣ、ＴｉＣ析出物をピン止め粒子として活用する手段を採用しているものがある（例えば、特許文献１〜３）。
【０００５】
これらの鋼材の製造方法としては、圧延加熱温度を１２００℃以上の高温にして、析出物を一旦溶体化し、熱間圧延後に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ＡｌＮ、Ｎｂ（ＣＮ）、Ｔｉ（ＣＮ）等の析出物を多量微細分散させている。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１２００℃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。
【０００６】
また、圧延ままの状態でも焼ならし処理と同じような微細組織とし、しかも浸炭処理時にオーステナイト結晶粒の粗大化を防止できる肌焼鋼を低温加熱圧延で製造する方法が種々提案されている（例えば、特許文献４）。
【０００７】
【特許文献１】
特開昭５７−８９４２５号公報
【特許文献２】
特開平８−１９９３０３号公報
【特許文献３】
特開平９−５９７４５号公報
【特許文献４】
特開昭５８−１１３３１８号公報
【０００８】
【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。
【０００９】
【課題を解決しようとする手段】
本発明者は、肌焼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１２００℃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１２００℃以上の高温として、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で分塊圧延した鋼片を低温加熱すると共に、圧延の全過程において一定の温度以上に維持する条件で熱間圧延すれば、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。
【００１０】
本発明の要旨は、次のとおりである。
【００１１】
（１）質量％で、
Ｃ：０．１〜０．４％、
Ｓｉ：０．０１〜１．３％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．０２〜０．１％、
Ｎ：０．００６〜０．０２５％
を含有し、さらに、
Ｃｒ：０．４〜１．８％、
Ｍｏ：０．０２〜１．０％、
Ｎｉ：０．１〜３．５％、
Ｖ：０．５％以下、
Ｔｉ：０．１％以下
の１種または２種以上を含有し、
Ｐ：０．０２５％以下、
Ｏ：０．００２５％以下
に制限し、残部が鉄及び不可避的不純物からなり、熱間圧延後のＡｌＮの析出量が０．０２％以上であり、マトリックス中に直径０．１μｍ以下のＡｌＮの析出物を５個／μｍ²以上を有し、硬さ指数Ｈを下記で定義すると、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％
【００１２】
（２）質量％で、
Ｃ：０．１〜０．４％、
Ｓｉ：０．０１〜１．３％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．０２〜０．１％、
Ｎｂ：０．００５〜０．０５％、
Ｎ：０．００６〜０．０２５％
を含有し、さらに、
Ｃｒ：０．４〜１．８％、
Ｍｏ：０．０２〜１．０％、
Ｎｉ：０．１〜３．５％、
Ｖ：０．５％以下、
Ｔｉ：０．１％以下
の１種または２種以上を含有し、
Ｐ：０．０２５％以下、
Ｏ：０．００２５％以下
に制限し、残部が鉄及び不可避的不純物からなり、熱間圧延後のＡｌＮの析出量が０．０２％以上であり、マトリックス中に直径０．１μｍ以下のＡｌＮ、Ｎｂ（ＣＮ）の析出物を５個／μｍ²以上を有し、硬さ指数Ｈを下記で定義すると、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】
（３）質量％で、
Ｃ：０．１〜０．４％、
Ｓｉ：０．０１〜１．３％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．０２〜０．１％、
Ｎ：０．００６〜０．０２５％
を含有し、さらに、
Ｃｒ：０．４〜１．８％、
Ｍｏ：０．０２〜１．０％、
Ｎｉ：０．１〜３．５％、
Ｖ：０．５％以下、
Ｔｉ：０．１％以下
の１種または２種以上を含有し、
Ｐ：０．０２５％以下、
Ｏ：０．００２５％以下
に制限し、残部が鉄及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、かつＡｌＮの析出量が０．００５％以下である鋼片を用い、加熱温度を９００〜１０７０℃、粗圧延から仕上げ圧延前までの圧延温度を８００℃以上、仕上げ温度を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ＡｌＮの析出量が０．０２％以上とし、マトリックス中に直径０．１μｍ以下のＡｌＮの析出物を５個／μｍ²以上とし、硬さ指数Ｈを下記で定義すると、硬さがＨＶでＨ＋３０以下とし、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下とす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％
【００１４】
（４）質量％で、
Ｃ：０．１〜０．４％、
Ｓｉ：０．０１〜１．３％、
Ｍｎ：０．３〜１．８％、
Ｓ：０．００１〜０．１５％、
Ａｌ：０．０２〜０．１％、
Ｎｂ：０．００５〜０．０５％、
Ｎ：０．００６〜０．０２５％
を含有し、さらに、
Ｃｒ：０．４〜１．８％、
Ｍｏ：０．０２〜１．０％、
Ｎｉ：０．１〜３．５％、
Ｖ：０．５％以下、
Ｔｉ：０．１％以下
の１種または２種以上を含有し、
Ｐ：０．０２５％以下、
Ｏ：０．００２５％以下
に制限し、残部が鉄及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、かつＡｌＮの析出量が０．００５％以下である鋼片を用い、加熱温度を９００〜１０７０℃、粗圧延から仕上げ圧延前までの圧延温度を８００℃以上、仕上げ温度を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ＡｌＮ、Ｎｂ（ＣＮ）の析出物を５個／μｍ²以上とし、硬さ指数Ｈを下記で定義すると、硬さがＨＶでＨ＋３０以下とし、ＪＩＳＧ０５５８で規定する脱炭深さ：ＤＭ−Ｔ０．２ｍｍ以下とす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鋼材の製造方法。
【００１５】
【発明の実施の形態】
以下、本発明について詳細に説明する。
【００１６】
本発明では、ＨＣＲ（ＨｏｔＣｈａｒｇｅＲｏｌｌｉｎｇ）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、ＡｌＮ、Ｎｂ（ＣＮ）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、ＪＩＳＧ０５５８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ともに軟質化により冷間加工性が向上できる。
【００１７】
まず、成分の限定理由について説明する。
【００１８】
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、０．１〜０．４％の範囲内にする必要がある。
【００１９】
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、１．３％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０．０１〜１．３％の範囲内にする必要がある。好適範囲は０．０２〜０．６％である。特に冷間加工性を重視する場合の好適範囲は、０．０２〜０．３％である。
【００２０】
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、０．３〜１．８％の範囲内にする必要がある。好適範囲は０．６〜１．０％である。
【００２１】
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。好適範囲は０．００５〜０．０４％である。
【００２２】
Ａｌは脱酸剤としての効果を有するとともに、鋼中にＡｌＮとして析出し、その後の浸炭処理においてオーステナイト結晶粒の粗大化を抑制する効果がある。Ａｌが０．０２％未満では、その効果が不十分であるが、０．１％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０．０２〜０．１％の範囲内にする必要がある。好適範囲は０．０２５〜０．０６％である。
【００２３】
また、ＮはＡｌＮを析出させるために必要な元素であり、そのためには０．００６％以上が必要である。一方、０．０２５％を超えると、析出するＡｌＮの凝集がその核発生よりも優先するので、粗大なＡｌＮが析出してしまい、かえってオーステナイト結晶粒の粗大化抑制効果が低下する。また、鋼の清浄度を害し、ブローホール発生の原因ともなる。したがって、Ｎは０．００６〜０．０２５％とした。
【００２４】
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。０．００５％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０．００５〜０．０５％の範囲内にする必要がある。好適範囲は、０．００５〜０．０３％である。
【００２５】
次に、本発明では、Ｃｒ、Ｍｏ、Ｎｉ、Ｖ、Ｔｉの１種または２種以上を選択成分として含有する。
【００２６】
即ち、Ｃｒ、Ｍｏ、Ｎｉは鋼に強度、焼入れ性を与えるのに有効な元素であるが、それぞれ、０．４％未満、０．０２％未満、０．１％未満ではその効果は不十分であり、一方、添加量が多すぎると鋼の硬さの上昇を招き冷間鍛造性が劣化するので、それぞれの上限を１．８％、１．０％、３．５％とした。したがって、Ｃｒ：０．４〜１．８％、Ｍｏ：０．０２〜１．０％、Ｎｉ：０．１〜３．５％とした。
【００２７】
また、Ｖ、Ｔｉは、鋼に強度を与えるとともに鋼中のＣ、Ｎと結びついてＶ（ＣＮ）、Ｔｉ（ＣＮ）を形成し、結晶粒の粗大化抑制に有効な元素であるが、それぞれ０．５％、０．１％を超えると、鋼が硬くなって冷間鍛造性を劣化させるので、Ｖ：０．５％、Ｔｉ：０．１％以下とした。
【００２８】
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。したがってその含有量を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。
【００２９】
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、Ａｌの析出物、Ｎｂの析出物の析出サイトとなり、熱間加工時にＡｌの析出物、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、Ｏ量はできるだけ低減することが望ましい。特に、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。
【００３０】
次に本発明では、熱間圧延後のＡｌＮの析出量が０．０２％以上であり、マトリックス中に直径０．１μｍ以下のＡｌＮの析出物を５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。
【００３１】
図１は、９７０℃で１時間の再加熱を行った際の結晶粒の粗大化率（結晶粒度Ｎｏ．４以下の粗大結晶粒の占める面積率）とＡｌＮ析出量との関係を示す図である。
【００３２】
図１に示すように、ＡｌＮ析出量が０．０２％以上となると結晶粒の粗大化が抑制できる。したがって、本発明ではＡｌＮ析出量を０．０２％以上とした。
【００３３】
また、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。
【００３４】
図２は、１μｍ²面積中の直径０．１μｍ以下のＡｌＮ析出物個数と粗大粒発生温度との関係を示す図である。なお、図２では圧下率５０％の据え込みを行った後、各温度で５時間浸炭でシュミレーションを行った結果を示している。図２から明らかなように、直径０．１μｍ以下のＡｌＮの微細析出物をその合計で５個／μｍ²以上分散させると実用上の浸炭加熱温度９７０℃以上の温度域において結晶粒の粗大化が生じず、優れた結晶粒粗大化防止特性が得られる。したがって、本発明では、マトリックス中に直径０．１μｍ以下のＡｌＮ析出物を５個／μｍ²以上分散させることとした。なお、好適範囲は１０個／μｍ²以上である。
【００３５】
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。
【００３６】
Ｈ＝２７３．５Ｃ％＋３９．１Ｓｉ％＋５４．７Ｍｎ％＋３０．４Ｃｒ％＋１３６．７Ｍｏ％＋１８．２Ｎｉ％
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、ＡｌＮ析出物を浸炭時に微細分散させることを特徴としている。本発明では、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ることを特徴としているが、分塊圧延後の鋼片中には、ＡｌＮの大部分は固溶状態にある。この鋼片を棒鋼線材圧延に際して低温加熱圧延すると、加熱過程で、オーステナイト中でＡｌＮが微細析出物する。この棒鋼線材圧延の加熱時にＡｌＮの一部が固溶し、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ＡｌＮやＮｂ（ＣＮ）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、Ｎｂの炭窒化物等が相界面析出し、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Ａｌ、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Ａｌ、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（Ａｌ、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Ａｌ、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Ａｌ、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、Ａｌ、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Ａｌ、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Ａｌ、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。
【００３７】
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−２５〜Ｈ＋２５の範囲である。
【００３８】
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。
【００３９】
次に本発明では、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。本発明者らは、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知見した。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。
【００４０】
【表１】

次に熱間圧延条件について説明する。
【００４１】
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、ＨＣＲ分塊圧延工程を経て、ＡｌＮの析出量が０．００５％以下である鋼片を用い、線材または棒鋼に低温加熱圧延して圧延素材とする。
【００４２】
加熱温度は９００〜１０７０℃のＡｒ₃点直上の温度として、熱間圧延途中の温度を８００℃以上に維持したまま粗圧延−仕上げ圧延前までの圧延を行い、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圧延し、熱間圧延後のＡｌＮの析出量が０．０２％以上で、マトリックス中に直径０．１μｍ以下のＡｌＮの析出物、または、ＡｌＮとＮｂ（ＣＮ）の析出物を５個／μｍ²以上とする。
【００４３】
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製造した鋼片を棒鋼線材圧延に際しての加熱過程で、オーステナイト中でＡｌＮが微細析出物する。加熱温度を９００〜１０７０℃のＡｒ₃点直上の温度とするのは、上記の微細なＡｌＮ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するので好ましくない。微細ＡｌＮ析出物をその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。
【００４４】
次に、粗圧延圧延から仕上げ圧延前までの圧延温度を８００℃以上、仕上げ温度を８００〜９７０℃とするのは次の理由による。熱間圧延途中の表面温度を８００℃未満、また仕上げ温度が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。上記の条件において、圧延材のフェライト脱炭が進行するのは、この温度域では、表層で未再結晶域圧延となり、フェライト変態が促進されること、あるいは圧延中に一部で歪み誘起変態が起こっていることが原因と考えられる。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、粗圧延圧延から仕上げ圧延前までの圧延温度を８００℃以上、仕上げ温度を８００〜９７０℃とする。好適温度は、粗圧延圧延から仕上げ圧延前までの圧延温度を８５０℃以上、仕上げ温度を８５０〜９６０℃である。
【００４５】
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。
【００４６】
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。
【００４７】
【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。
【００４８】
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Ｓ、Ｔ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）。
【００４９】
【表２】

【００５０】
続いて、熱間圧延により、直径３２ｍｍの棒鋼を製造した。熱間圧延条件を表３に示す。熱間圧延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。
【００５１】
熱間圧延後の棒鋼のＡｌ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０．１μｍ以下のＡｌの析出物、Ｎｂの析出物、ＡｌとＮｂの複合組成からなる析出物、Ｖ、Ｔｉ添加鋼についてはこれらの析出物の数を数え、１平方μｍ当たりの数に換算した。
【００５２】
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。
【００５３】
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ＪＩＳＧ０５５１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。
【００５４】
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入れを行い、γ粒度を測定した。
【００５５】
これらの調査結果を熱間圧延条件とあわせて表３に示す。
【００５６】
比較例２５、２６、２７は鋼水準Ａ、Ｓ、Ｔを従来の製造条件で製造した鋼材の特性であるが、本発明例の冷間変形抵抗は、各々同一成分系（Ｃｒ系、Ｍｏ系）の１、２、９、１０、１１と比較すると、各々比較例２５、２６、２７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。比較例２５、２６、２７に比較して顕著に優れている。
【００５７】
次に、表３において、比較例１６はＡｌ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１７はＮ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１８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。
【００５８】
比較例１９、２０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。
【００５９】
比較例２１は熱間圧延前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冷間加工性も劣り、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。
【００６０】
比較例２２は熱間圧延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２に比較して冷間加工性は劣る。比較例２３は粗圧延から仕上げ圧延前までの圧延温度が本願規定の範囲を下回り、さらに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２４は熱間圧延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。
【００６１】
【表３】

【００６２】
【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】９７０℃で１時間の再加熱を行った際の結晶粒の粗大化率（結晶粒度Ｎｏ．４以下の粗大結晶粒の占める面積率）とＡｌＮ析出量との関係を示す図である。
【図２】１μｍ²面積中の直径０．１μｍ以下のＡｌＮ析出物個数と粗大粒発生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.
[0003]
Case-hardened steel is liable to cause a phenomenon that some austenite crystal grains become coarse during carburizing heating. Parts with such coarse grains generate heat treatment distortion after carburizing and quenching. For example, gears and CVJ parts with large carburizing distortion cause noise and vibration.
[0004]
Conventionally, in order to prevent coarsening of austenite crystal grains in carburizing treatment, the presence form of AlN in the steel material after hot rolling, the amount of precipitation of fine precipitates such as Nb (CN), Ti (CN), etc. Is used to control NbC and TiC precipitates as pinning particles (for example, Patent Documents 1 to 3).
[0005]
As a manufacturing method of these steel materials, the heating temperature after rolling is set to a high temperature of 1200 ° C. or higher, the precipitate is once solutionized, and after the hot rolling, the precipitation temperature range of the precipitate is gradually cooled, so that A large amount of precipitates such as AlN, Nb (CN), and Ti (CN) are finely dispersed in the steel material. However, even in these methods, the prevention of coarse grains during carburizing may not always be sufficient, and when the heating temperature for rolling is 1200 ° C. or higher, the rolling becomes hard and cold workability becomes insufficient, In addition, there is a problem that total decarburization becomes prominent.
[0006]
In addition, various methods have been proposed for producing a case-hardened steel by low-temperature hot rolling that has the same microstructure as that of the normalizing process even in the rolled state, and that can prevent austenite crystal grains from becoming coarse during the carburizing process ( For example, Patent Document 4).
[0007]
[Patent Document 1]
JP-A-57-89425 [Patent Document 2]
JP-A-8-199303 [Patent Document 3]
JP-A-9-59745 [Patent Document 4]
Japanese Patent Laid-Open No. 58-113318
[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.
[0009]
[Means to solve the problem]
The present inventor has intensively studied to improve the cold workability of case-hardened steel and to improve the coarse grain prevention characteristics during carburizing. As a result, when the rolling heating temperature is set to a high temperature of 1200 ° C. or higher as in the conventional case, when the total decarburization becomes significant, coarse particles are generated from the decarburized part at the time of subsequent carburization, and thus the coarse particle prevention characteristics are insufficient. I found out. Furthermore, the steel sheet that has been subjected to HCR split rolling with HCR without cooling to the A3 point temperature or less after casting, even if the rolling heating temperature is set to a high temperature of 1200 ° C. or higher as in the prior art and the precipitate is not once solutionized, is heated at a low temperature. In addition, if the hot rolling is performed under conditions that maintain a certain temperature or higher in the entire rolling process, it is possible to reduce the total decarburization, thereby preventing the formation of coarse grains and reducing the amount of precipitation strengthening. The present invention was completed by finding that cold workability can be improved by softening by reduction.
[0010]
The gist of the present invention is as follows.
[0011]
(1) In mass%,
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: Restricted to 0.0025% or less, the balance is made of iron and inevitable impurities, the precipitation amount of AlN after hot rolling is 0.02% or more, and AlN having a diameter of 0.1 μm or less in the matrix When the number of precipitates is 5 / μm ² or more and the hardness index H is defined below, it is H + 30 or less in HV, and the decarburization depth specified in JIS G 0558: DM-T 0.2 mm or less. A steel for case hardening with excellent cold workability and characteristics of preventing coarse grains during carburizing.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0012]
(2) By mass%
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
Nb: 0.005 to 0.05%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: limited to 0.0025% or less, the balance being iron and inevitable impurities, the precipitation amount of AlN after hot rolling is 0.02% or more, AlN having a diameter of 0.1 μm or less in the matrix, Decarburization depth defined by JIS G 0558: DM having Nb (CN) precipitates of 5 / μm ² or more and a hardness index H defined as follows: -A steel for skin hardening excellent in cold workability and coarse grain prevention characteristics during carburizing, characterized by being T 0.2 mm or less.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0013]
(3) In mass%,
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: The steel is limited to 0.0025% or less, and the balance is manufactured by a step of performing a piece rolling without cooling the steel composed of iron and inevitable impurities to A3 point temperature or less after casting, and the precipitation amount of AlN is A steel slab of 0.005% or less is used, the heating temperature is 900 to 1070 ° C., the rolling temperature from rough rolling to finish rolling is 800 ° C. or higher, the finishing temperature is 800 to 970 ° C., and hot rolling is followed by 800. The steel is hot-rolled to a wire rod or steel bar under the condition that the temperature range of ˜500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less, the precipitation amount of AlN after hot rolling is 0.02% or more, Deposition of AlN having a diameter of 0.1 μm or less is 5 / μm ² or more, and the hardness index H is defined as follows, the hardness is HV and H + 30 or less, and the decarburization depth specified in JIS G 0558: DM -T0.2mm or less Cold workability and a manufacturing method excellent hardened steel material in preventing coarse grains characteristic during carburization, characterized in that a.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
[0014]
(4) By mass%
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
Nb: 0.005 to 0.05%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: The steel is limited to 0.0025% or less, and the balance is manufactured by a step of performing a piece rolling without cooling the steel composed of iron and inevitable impurities to A3 point temperature or less after casting, and the precipitation amount of AlN is A steel slab of 0.005% or less is used, the heating temperature is 900 to 1070 ° C., the rolling temperature from rough rolling to finish rolling is 800 ° C. or higher, the finishing temperature is 800 to 970 ° C., and hot rolling is followed by 800. It is hot-rolled to a wire rod or steel bar under conditions where the temperature range of ˜500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less, and 5 Al / Nb (CN) precipitates after hot rolling are 5 pieces / μm ^2. When the hardness index H is defined below, the cold workability is characterized in that the hardness is H + 30 or less and the decarburization depth specified in JIS G 0558 is DM-T 0.2 mm or less. And prevention of coarse grains during carburizing Method of manufacturing the excellent case hardened steel.
[0015]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be described in detail.
[0016]
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 being cast, and cold rolling at the time of steel bar wire rolling, AlN, Nb ( CN) fine precipitates are kept as they are. 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.
[0017]
First, the reasons for limiting the components will be described.
[0018]
C is an element effective for giving steel the necessary strength, 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%.
[0019]
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 1.3%, 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 1.3%. The preferred range is 0.02 to 0.6%. A preferable range when particularly emphasizing cold workability is 0.02 to 0.3%.
[0020]
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 0.6-1.0%.
[0021]
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%. The preferred range is 0.005 to 0.04%.
[0022]
Al has an effect as a deoxidizing agent, and also has an effect of precipitating as AlN in the steel and suppressing coarsening of austenite crystal grains in the subsequent carburizing treatment. If Al is less than 0.02%, the effect is insufficient, but if it exceeds 0.1%, AlN remains without solution during rolling and heating, and becomes a precipitation site for Nb and Ti precipitates. It inhibits fine dispersion of precipitates and promotes coarsening of crystal grains. For the above reasons, the content needs to be in the range of 0.02 to 0.1%. The preferred range is 0.025 to 0.06%.
[0023]
N is an element necessary for precipitating AlN, and 0.006% or more is necessary for this purpose. On the other hand, if it exceeds 0.025%, the aggregation of precipitated AlN has priority over the generation of nuclei, so that coarse AlN is precipitated, and on the contrary, the effect of suppressing the coarsening of austenite crystal grains decreases. Moreover, it impairs the cleanliness of the steel and causes blowholes. Therefore, N is set to 0.006 to 0.025%.
[0024]
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 0.005%, 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 0.005 to 0.05%. The preferred range is 0.005 to 0.03%.
[0025]
Next, in this invention, 1 type (s) or 2 or more types of Cr, Mo, Ni, V, and Ti are contained as a selection component.
[0026]
That is, Cr, Mo, and Ni are effective elements for imparting strength and hardenability to steel, but the effects are insufficient when less than 0.4%, less than 0.02%, and less than 0.1%, respectively. On the other hand, if the addition amount is too large, the hardness of the steel is increased and the cold forgeability is deteriorated. Therefore, the upper limits are set to 1.8%, 1.0%, and 3.5%, respectively. Therefore, Cr: 0.4-1.8%, Mo: 0.02-1.0%, Ni: 0.1-3.5%.
[0027]
V and Ti are elements that give strength to the steel and are combined with C and N in the steel to form V (CN) and Ti (CN), and are effective in suppressing the coarsening of crystal grains. If it exceeds 0.5% and 0.1%, the steel becomes hard and the cold forgeability deteriorates, so V: 0.5% and Ti: 0.1% or less.
[0028]
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.
[0029]
O forms oxide inclusions such as Al ₂ O _{3 in} the steel. If a large amount of oxide inclusions are present in the steel, it becomes a precipitation site for Al precipitates and Nb precipitates, during which Al precipitates and Nb 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. In particular, if 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.
[0030]
Next, in the present invention, the precipitation amount of AlN after hot rolling is 0.02% or more, and the matrix has 5 precipitations / μm ² of AlN having a diameter of 0.1 μm or less. The reason for limiting to is described below.
[0031]
1 is a diagram showing the relationship between the coarsening rate of crystal grains (area ratio occupied by coarse crystal grains having a grain size of No. 4 or less) and the amount of precipitated AlN when reheating is performed at 970 ° C. for 1 hour. is there.
[0032]
As shown in FIG. 1, when the AlN precipitation amount is 0.02% or more, coarsening of crystal grains can be suppressed. Therefore, in the present invention, the AlN precipitation amount is set to 0.02% or more.
[0033]
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 particle diameter and the larger the amount, the more pinned particles This is preferable because the number of
[0034]
FIG. 2 is a graph showing the relationship between the number of AlN precipitates having a diameter of 0.1 μm or less in a 1 μm ² area and the coarse grain generation temperature. FIG. 2 shows the result of simulation by carburizing for 5 hours at each temperature after upsetting with a reduction ratio of 50%. As is apparent from FIG. 2, when AlN fine precipitates having a diameter of 0.1 μm or less are dispersed in a total of 5 particles / μm ² or more, crystal grains become coarse in a practical carburizing heating temperature of 970 ° C. or more. Does not occur, and excellent crystal grain coarsening prevention characteristics can be obtained. Accordingly, in the present invention, 5 / μm ² or more AlN precipitates having a diameter of 0.1 μm or less are dispersed in the matrix. The preferred range is 10 / μm ² or more.
[0035]
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.
[0036]
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%
Claim 1 or 2 of the present invention is characterized in that AlN precipitates are finely dispersed during carburizing in order to prevent coarse grains during carburizing. In the present invention, it is characterized by using a steel slab produced by performing the partial rolling without cooling below the A3 point temperature after casting, but in the steel slab after the partial rolling, most of AlN is It is in a solid solution state. When this steel slab is rolled at a low temperature during the rolling of the bar steel wire rod, AlN finely precipitates in the austenite during the heating process. A part of AlN dissolves during heating of the steel bar wire rolling, and the precipitate grows Ostwald. When the precipitate grows through Ostwald, the coarse grain prevention characteristics during carburization deteriorate. Moreover, when AlN and Nb (CN) are partly dissolved during heating in the steel bar wire rolling, Nb carbonitrides and the like are precipitated at the interface between the austenite and ferrite during the cooling process after the steel bar wire rolling. Hardness increases by precipitation hardening. In other words, the hardness of the rolled steel bar reflects the amount of Al and Nb carbonitrides that are dissolved during heating of the rolled steel bar. The harder the rolled steel bar is, the harder it is. The amount of carbonitrides of Al and Nb that are solid-dissolved during heating in the rolling of the steel bar wire 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 Al and Nb), precipitation hardening due to Al and Nb in the cooling process during rolling of the steel bar is reduced. This makes it possible to finely disperse Al and Nb precipitates during carburizing, and to prevent coarse particles during 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 Al 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 Al 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 Al 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.
[0037]
When the hardness of the hot-worked material exceeds HV at H + 30, the hardness of the hot-worked material becomes hard and cold workability deteriorates, so the hardness of the hot-worked material is limited to a range of H + 30 or less in HV. . The preferred range is H-25 to H + 25.
[0038]
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.
[0039]
Next, in the present invention, the upper limit of the decarburization depth is defined for the purpose of preventing coarse grains. 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 inventors have found for the first time that when the decarburization depth exceeds DM-T 0.2 mm, coarse grains are likely to be generated during carburization. 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.
[0040]
[Table 1]

Next, hot rolling conditions will be described.
[0041]
The steel of the present invention is melted by a usual method such as a converter, an electric furnace, etc., the components are adjusted, and after casting, without cooling to A3 point or less, through the HCR block rolling process, the AlN A steel slab having a precipitation amount of 0.005% or less is used, and a wire rod or a steel bar is subjected to low temperature heating rolling to obtain a rolled material.
[0042]
The heating temperature is a temperature just above the Ar ₃ point of 900 to 1070 ° C., and the rolling is performed before rough rolling-finishing rolling while maintaining the temperature during hot rolling at 800 ° C. or higher, and the finishing temperature is 800 to 970 ° C. Hot rolling is performed. Subsequent to hot rolling, a wire or bar steel is hot-rolled under the condition of gradually cooling a temperature range of 800 to 500 ° C. at a cooling rate of 1 ° C./second or less, and the precipitation amount of AlN after hot rolling is 0.02 %, And the number of AlN precipitates having a diameter of 0.1 μm or less or AlN and Nb (CN) precipitates in the matrix is 5 / μm ² or more.
[0043]
In the austenite, AlN is finely precipitated in the austenite during the heating process in the steel bar wire rolling of the steel slab produced by performing the batch rolling without cooling below the A3 point temperature after casting. The reason why the heating temperature is set to a temperature immediately above the Ar ₃ point of 900 to 1070 ° C. is to prevent the fine AlN precipitates from being dissolved in the matrix. This is not preferable, and if it exceeds 1070 ° C., the precipitate is dissolved in the matrix, which is not preferable. By maintaining the fine AlN precipitate as it is, the heating temperature was set to 900 to 1070 ° C. in order to suppress the generation of coarse particles during carburization.
[0044]
Next, the rolling temperature from rough rolling to finish rolling is set to 800 ° C. or higher, and the finishing temperature is set to 800 to 970 ° C. for the following reason. If the surface temperature during hot rolling is less than 800 ° C and the finish temperature is less than 800 ° C, ferrite decarburization of the rolled material proceeds. As a result, total decarburization becomes remarkable, and coarse grains are generated during carburizing. It becomes easy to do. Under the above-mentioned conditions, the ferrite decarburization of the rolled material proceeds in this temperature range, in which the surface layer is non-recrystallized zone rolling, and the ferrite transformation is promoted, or some strain-induced transformation occurs during rolling. The cause is thought to be happening. 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 these reasons, the rolling temperature from rough rolling to finish rolling is set to 800 ° C. or higher, and the finishing temperature is set to 800 to 970 ° C. The preferable temperature is 850 ° C. or higher for the rolling temperature from rough rolling to finish rolling, and the finishing temperature is 850 to 960 ° C.
[0045]
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.
[0046]
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.
[0047]
【Example】
Hereinafter, the effects of the present invention will be described more specifically by way of examples.
[0048]
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 S and T, after continuous casting, the steel was once cooled to room temperature, and then heated again to the A3 point or higher to perform a batch rolling to obtain a 162 mm square steel slab (rolling material). Condition II).
[0049]
[Table 2]

[0050]
Subsequently, a steel bar having a diameter of 32 mm was manufactured by hot rolling. Table 3 shows the hot rolling conditions. The cooling rate after hot rolling was adjusted using a slow cooling cover installed on the cooling bed.
[0051]
In order to investigate the dispersion state of Al precipitates and Nb precipitates in the steel bar after hot rolling, 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 about 30,000 times, and about 20 fields of view are observed. About the precipitate of Al having a diameter of 0.1 μm or less in one field of view, the precipitate of Nb, the precipitate composed of a composite composition of Al and Nb, V, and Ti-added steel The number of these precipitates was counted and converted into the number per square μm.
[0052]
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.
[0053]
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.
[0054]
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.
[0055]
These investigation results are shown in Table 3 together with hot rolling conditions.
[0056]
Comparative Examples 25, 26, and 27 are the characteristics of steel materials manufactured with steel levels A, S, and T under conventional manufacturing conditions. The cold deformation resistances of the examples of the present invention are the same component system (Cr system, Mo system). ), 1, 2, 9, 10, and 11 are significantly smaller than those of Comparative Examples 25, 26, and 27, respectively, 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. Compared with Comparative Examples 25, 26, and 27, it is remarkably superior.
[0057]
Next, in Table 3, Comparative Example 16 is a case where the Al content falls below the range specified in the present application, and Comparative Example 17 is a case where the N content falls below the range specified in the present application. Is a case where the content of O exceeds the range specified in the present application, and the coarse grain prevention characteristics are all inferior.
[0058]
Comparative Examples 19 and 20 are different from the present invention in the method of manufacturing the steel slab, and after casting, the steel was once cooled to the A3 point temperature or lower and then subjected to the lump rolling method, both having coarse grain prevention characteristics. Is inferior.
[0059]
Comparative Example 21 is a case where the heating temperature before hot rolling exceeds the range specified in the present application, the number of precipitates is less than the range of the present invention, and the hardness after rolling exceeds the range specified in the present application, and cold workability. In addition, the total decarburization depth exceeds the range of the present invention, and the coarse grain prevention property is inferior.
[0060]
Comparative Example 22 is a case where the hot rolling finishing temperature exceeds the range specified in the present application, and the cold workability is inferior to that of Example 2 of the present invention. Comparative Example 23 is a case where the rolling temperature from rough rolling to finish rolling is lower than the range specified in the present application, and the finishing temperature is lower than the range specified in the present application, and the total decarburization depth exceeds the range of the present invention, Coarse grain prevention properties are inferior. The comparative example 24 is a case where the cooling rate after hot rolling exceeds the range specified in the present application, and the coarse grain prevention property is inferior and the cold workability is also inferior.
[0061]
[Table 3]

[0062]
【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. The generation of coarse particles can be stably suppressed, thereby preventing the occurrence of distortion and bending. Therefore, it has become 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]
1 is a graph showing the relationship between the coarsening rate of crystal grains (area ratio occupied by coarse crystal grains having a grain size of No. 4 or less) and the amount of precipitated AlN when reheating is performed at 970 ° C. for 1 hour. is there.
FIG. 2 is a diagram showing the relationship between the number of AlN precipitates having a diameter of 0.1 μm or less in 1 μm ² area and the generation temperature of coarse particles.

Claims

% By mass
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: Restricted to 0.0025% or less, the balance is made of iron and inevitable impurities, the precipitation amount of AlN after hot rolling is 0.02% or more, and AlN having a diameter of 0.1 μm or less in the matrix When the number of precipitates is 5 / μm ² or more and the hardness index H is defined below, it is H + 30 or less in HV, and the decarburization depth specified in JIS G 0558: DM-T 0.2 mm or less. A steel for case hardening with excellent cold workability and characteristics of preventing coarse grains during carburizing.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%

% By mass
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
Nb: 0.005 to 0.05%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: limited to 0.0025% or less, the balance being iron and inevitable impurities, the precipitation amount of AlN after hot rolling is 0.02% or more, AlN having a diameter of 0.1 μm or less in the matrix, Decarburization depth defined by JIS G 0558: DM having Nb (CN) precipitates of 5 / μm ² or more and a hardness index H defined as follows: -A steel for skin hardening excellent in cold workability and coarse grain prevention characteristics during carburizing, characterized by being T 0.2 mm or less.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%

% By mass
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: The steel is limited to 0.0025% or less, and the balance is manufactured by a step of performing a piece rolling without cooling the steel composed of iron and inevitable impurities to A3 point temperature or less after casting, and the precipitation amount of AlN is A steel slab of 0.005% or less is used, the heating temperature is 900 to 1070 ° C., the rolling temperature from rough rolling to finish rolling is 800 ° C. or higher, the finishing temperature is 800 to 970 ° C., and hot rolling is followed by 800. The steel is hot-rolled to a wire rod or steel bar under the condition that the temperature range of ˜500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less, the precipitation amount of AlN after hot rolling is 0.02% or more, Deposition of AlN having a diameter of 0.1 μm or less is 5 / μm ² or more, and the hardness index H is defined as follows, the hardness is HV and H + 30 or less, and the decarburization depth specified in JIS G 0558: DM -T0.2mm or less Cold workability and a manufacturing method excellent hardened steel material in preventing coarse grains characteristic during carburization, characterized in that a.
H = 273.5C% + 39.1Si + 54.7Mn% + 30.4Cr% + 136.7Mo% + 18.2Ni%

% By mass
C: 0.1-0.4%
Si: 0.01 to 1.3%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: 0.02 to 0.1%,
Nb: 0.005 to 0.05%,
N: 0.006 to 0.025%
In addition,
Cr: 0.4 to 1.8%,
Mo: 0.02 to 1.0%,
Ni: 0.1 to 3.5%
V: 0.5% or less,
Ti: 0.1% or less containing one or more,
P: 0.025% or less,
O: The steel is limited to 0.0025% or less, and the balance is manufactured by a step of performing a piece rolling without cooling the steel composed of iron and inevitable impurities to A3 point temperature or less after casting, and the precipitation amount of AlN is A steel slab of 0.005% or less is used, the heating temperature is 900 to 1070 ° C., the rolling temperature from rough rolling to finish rolling is 800 ° C. or higher, the finishing temperature is 800 to 970 ° C., and hot rolling is followed by 800. It is hot-rolled to a wire rod or steel bar under conditions where the temperature range of ˜500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less, and 5 Al / Nb (CN) precipitates after hot rolling are 5 pieces / μm ^2. When the hardness index H is defined below, the cold workability is characterized in that the hardness is H + 30 or less and the decarburization depth specified in JIS G 0558 is DM-T 0.2 mm or less. And prevention of coarse grains during carburizing Method of manufacturing the excellent case hardened steel.