JP4121416B2

JP4121416B2 - Non-tempered hot forged parts for machine structure and manufacturing method thereof

Info

Publication number: JP4121416B2
Application number: JP2003120405A
Authority: JP
Inventors: 哲夫白神; 和明福岡; 邦和冨田; 義正船川; 毅塩崎
Original assignee: JFE Steel Corp; JFE Bars and Shapes Corp
Current assignee: JFE Steel Corp; JFE Bars and Shapes Corp
Priority date: 2002-04-26
Filing date: 2003-04-24
Publication date: 2008-07-23
Anticipated expiration: 2023-04-24
Also published as: JP2004003011A

Description

【０００１】
【発明の属する技術分野】
本発明は機械構造用非調質型熱間鍛造部品およびその製造方法に関し、特に熱間鍛造後非調質であっても、引張強さ７００ＭＰａ以上、且つ降伏比が０．８５以上と高く、疲労特性および機械加工性に優れるものに関する。
【０００２】
【従来の技術】
自動車や建設機械に用いられる構造用部品として機械構造用炭素鋼や機械構造用合金鋼を部品形状に熱間鍛造後、焼入れ焼戻する調質部品や、焼入れ焼戻しを省略した非調質部品が用いられている。
【０００３】
非調質部品の場合、フェライト・パーライト二相組織のため、引張強さは８００ＭＰａ以下と低いものが多く、強度向上が課題とされてきた。
【０００４】
しかしながら、非調質部品の強度向上において、降伏強度を高くし、疲労強度を調質部品並とした場合、鋼が硬化し、靭性や機械加工性が劣化することが懸念されるため、降伏比を高くする、すなわち、降伏強度は高くするが引張強さは高くしないことが必要である。
【０００５】
特開平１１−２９８４２号公報はフェライト・パーライトの二相組織において引張強さ、疲労強度を向上させる技術に関するものであるが、鋼組成においてＣ量を増大させるため、靭性の低下が懸念される。
【０００６】
特開平７−１０９５１８号公報は非調質鋼の強度、降伏比を向上させるため、低Ｃ鋼をフェライト・ベイナイトの二相組織とし、鍛造後に時効熱処理することを提案している。本技術によれば引張強さは向上するものの降伏比は０．８以下と低い。
【０００７】
特開２０００−１２９３９３号公報は熱間鍛造用非調質鋼の強度、靭性を向上させるため鋼組成を低炭素−高合金系とし焼入れ性を向上させマルテンサイト・ベイナイト組織とすることを提案している。しかしながら、フェライト・パーライト組織鋼より低い降伏比が記載され、疲労強度の低下が懸念される。
【０００８】
特開平７−１７３５３１号公報は非調質鋼の降伏比を向上させるため熱間鍛造後、ベイナイト・マルテンサイト組織とし、焼戻しを行うことを提案しているが、調質部品の降伏比には及ばない。
【０００９】
【発明が解決しようとする課題】
上述したように、調質部品に匹敵する靭性、降伏強度および降伏比の全てを満足する非調質部品用鋼材はいまだ開発されていないのが現状である。
【００１０】
そこで本発明では、これらの特性において調質部品に匹敵する性能を有し、被削性に優れた機械構造用非調質型熱間鍛造部品およびその製造方法を提供することを目的とする。
【００１１】
【課題を解決するための手段】
本発明者等は、非調質材で調質材に匹敵する強度、靭性が得られ、且つ被削性を損なわない鋼材についてその成分組成、製造条件の観点から種々検討し、鋼を微細析出物を利用した析出強化により強化した場合、降伏強度は上昇するが高降伏比となるため、鋼が過度に硬化し靭性や被削性を損なうことのないことを見出した。
【００１２】
本発明は以上の知見を基に更に検討を加えてなされたものである。すなわち、本発明は、
１．鋼組成が、質量％で、Ｃ≦０．１５％、Ｓｉ≦１％、Ｍｎ≦２％、Ｔｉ：０．０３〜０．３５％、Ｍｏ：０．０５〜０．８％、残部Ｆｅ及び不可避的不純物からなり、フェライト単相組織を有し、フェライト相中の、粒径が１０ｎｍ未満の微細析出物の占める割合が全析出物の９０％以上で、下記式（１）
０．５≦（Ｃ／１２）／｛（Ｔｉ／４８）+（Ｍｏ／９６）｝≦１．５ ---（１）
但し、上記式（１）において、各元素は、含有量（質量％）とする。
を満足することを特徴とする、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品。
【００１３】
２．前記微細析出物がＴｉ、Ｍｏの炭化物であることを特徴とする、請求項１記載の、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品。
【００１４】
３．鋼組成が、質量％で、Ｃ≦０．１５％、Ｓｉ≦１％、Ｍｎ≦２％、Ｔｉ：０．０３〜０．３５％、Ｍｏ：０．０５〜０．８％を含有し、更に、質量％で、Ｎｂ≦０．０８％、Ｖ ≦０．１５％、Ｗ≦１．５％の一種または二種以上を含有し、下記式（２）
０．５≦（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）
＋（Ｖ／５１）＋（Ｗ／１８４）｝≦１．５ ---（２）
但し、上記式（２）において、各元素は、含有量（質量％）とし、含まれないものは０とする。
を満足することを特徴とする、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品。
【００１５】
４．前記微細析出物がＴｉとＭｏとＮｂ、Ｖ、Ｗの内の少なくとも一種とを含む炭化物であることを特徴とする、請求項３記載の、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品。
【００２０】
５．鋼組成が請求項１または３の記載の鋼を１１００℃以上に加熱後、９００℃以上１２００℃以下の温度で部品形状に熱間鍛造し、その後の冷却において、５５０〜７００℃を０．５℃／ｓｅｃ以下の冷却速度で冷却することを特徴とする、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品の製造方法。
【００２１】
６．鋼組成が請求項１または３に記載の鋼を１１００℃以上に加熱後、９００℃以上１２００℃以下の温度で熱間鍛造し、その後の冷却過程において、５５０〜７００℃に１０分以上６０分以下で保持することを特徴とする、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品の製造方法。
【００２２】
７．鋼組成が請求項１または３に記載の鋼を１１００℃以上に加熱後、９００℃以上１２００℃以下の温度で熱間鍛造し、５５０〜７００℃に加熱して１０分以上６０分以下で保持することを特徴とする、部品の任意の位置から切り出した試験片の引張強さが７００ＭＰａ以上、降伏比が０．８５以上である機械構造用非調質型熱間鍛造部品の製造方法。
【００２３】
【発明の実施の形態】
本発明のミクロ組織、成分組成および製造条件について以下に詳細に説明する。
【００２４】
１．ミクロ組織
本発明に係る熱間鍛造部品はそのミクロ組織を、フェライト単相組織に粒径１０ｎｍ未満の微細析出物を分散析出させた組織とする。
フェライト単相組織とした場合、調質材に匹敵する靭性が得られ、該組織中に微細析出物を分散析出させた場合、調質材に匹敵する強度特性（高降伏強度、高降伏比）が得られる。微細析出物は熱間鍛造後の冷却速度の調整や、熱間鍛造後、析出処理により析出させる。
【００２５】
本発明では微細析出物は粒径１０ｎｍ未満とする。析出物の粒径が１０ｎｍ以上の場合、輸送機器、建機の機械構造用部品として必要な引張強さ７００ＭＰａ以上が得られない。
【００２６】
また、フェライト単相組織中に粒径１０ｎｍ未満の微細析出物を析出させた場合、降伏比が０．８５以上となり、降伏強度の上昇に対して引張強度の上昇、すなわち、鋼の硬化が小さく調質鋼に匹敵する被削性が得られる。母相、析出物のいずれかが本発明の規定外となった場合、降伏比は０．８５未満となる。
【００２７】
微細析出物の粒径は小さいほど強度向上に有効で、望ましくは５ｎｍ、更に望ましくは３ｎｍ以下とし、そのような微細析出物としてＴｉ、Ｍｏを複合含有した炭化物、またそれらに更にＮｂ、Ｖ、Ｗの一種または二種以上を含む炭化物が好ましい。
【００２８】
これらの微細析出物の分布形態は特に規定しないが、母相中に均一分散（分散析出）することが望ましい。
【００２９】
また、本発明において、微細析出物の大きさは、全析出物の９０％以上で満足すれば、目的とする引張強さ７００ＭＰａ以上が得られる。但し、１０ｎｍ以上の大きさの析出物は析出物形成元素を消費し、強度に悪影響をあたえるため、５０ｎｍ以下とすることが好ましい。
【００３０】
上述した析出物とは別に少量のＦｅ炭化物を含有しても本発明の効果は損なわれないが、平均粒径が１μｍ以上のＦｅ炭化物を多量に含むと靭性を阻害するため、本発明においては含有されるＦｅ炭化物の大きさ上限は１μｍ、含有率は全体の１％以下とすることが望ましい。
【００３１】
微細析出物の全析出物に占める割合は、以下の方法により求める。電子顕微鏡試料を、ツインジェット法を用いた電解研磨法で作成し、加速電圧２００ｋＶで観察する。その際、微細析出物が母相に対して計測可能なコントラストになるように母相の結晶方位を制御し、析出物の数え落としを最低限にするために焦点を正焦点からずらしたデフォーカス法で観察を行う。
【００３２】
また、析出物粒子の計測を行った領域の試料の厚さは電子エネルギー損失分光法を用いて、弾性散乱ピークと非弾性散乱ピーク強度を測定することで評価する。
【００３３】
この方法により、粒子数の計測と試料厚さの計測を同じ領域について実行することができる。粒子数および粒子径の測定は試料の０．５×０．５μｍの領域４箇所について行い、１μｍ²当たりに分布する析出物を粒径ごとの個数として算出する。
【００３４】
この値と試料厚さから、析出物の１μｍ³当たりに分布する粒子径ごとの個数を算出し、径が１０ｎｍ未満の析出物について、測定した全析出物に占める割合を算出する。
【００３５】
また、本発明においてフェライト単相組織とは、断面組織観察（２００倍の光学顕微鏡組織観察）でフェライト面積率９５％以上とし、好ましくは９８％以上とする。
【００３６】
２．成分組成
本発明鋼は上述したミクロ組織で目的とする性能が得られるが、以下の成分組成が好ましい。
【００３７】
Ｃ
Ｃは０．１５％を超えて含有すると微細析出物が粗大化し、強度が低下するため０．１５％以下とすることが好ましい。より好ましくは、０．０３％以上、０．１２％以下である。
【００３８】
Ｓｉ
Ｓｉは脱酸のために添加するが、１％を超えるとフェライトに固溶し、冷間加工時の変形抵抗が増大するため１％以下とする。より好ましくは、０．１５％以下である。
【００３９】
Ｍｎ
Ｍｎは強度向上に有効なため添加するが、２％を超えると冷間加工性を劣化させるので、２％以下とする。より好ましくは、０．５％以上、１．８％以下である。
【００４０】
Ｔｉ
ＴｉはＴｉ系炭化物、ＭｏとともにＴｉ−Ｍｏ系炭化物を含む析出物を微細に析出させ、強度を向上させるため添加する。引張強度７００ＭＰａ以上を確保するため０．０３％以上とし、一方、０．３５％を超えて添加すると析出物が粗大化し、強度が低下するため０．０３〜０．３５％とする。より好ましくは、０．０３〜０．２０％である。
【００４１】
Ｍｏ
ＭｏはＭｏ系炭化物、ＴｉとともにＴｉ−Ｍｏ系炭化物を含む析出物を微細に析出させ、強度を向上させるため添加する。引張強度７００ＭＰａ以上を確保するため０．０５％以上とし、一方、０．８％を超えて添加するとベイナイト等の低温変態相を形成し、微細析出物による析出強化が不足し、強度が低下するため０．０５〜０．８％とする。より好ましくは、０．１５〜０．４５％である。
【００４２】
Ｍｏは拡散速度が遅く、Ｔｉとともに析出する場合、析出物の成長速度が低下し、微細な析出物が得られやすい。
【００４３】
（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）｝
本パラメータは、析出物の大きさに影響を与えるもので、０．５以上、１．５以下とした場合、粒径１０ｎｍ未満の微細析出物の形成が容易となり好ましい。さらに好ましくは、０．７以上１．３以下である。
【００４４】
微細なＴｉ−Ｍｏ系炭化物では、炭化物中のＴｉ、Ｍｏは原子比で０．２≦Ｔｉ／Ｍｏ≦２．０、更に微細な炭化物では０．７≦Ｔｉ／Ｍｏ≦１．５であることが観察された。
【００４５】
更に、特性を向上させる場合、Ｎｂ、Ｖ、Ｗの一種または二種以上を添加することが好ましい。
【００４６】
Ｎｂ
ＮｂはＴｉとともに微細析出物を形成して強度上昇に寄与する。また、組織を微細化し、また結晶粒の整粒により延性を向上させる。０．０８％を超えると過度に微細化し、延性が低下するため０．０８％以下とする。より好ましくは０．０４％以下である。
【００４７】
Ｖ
ＶはＴｉと微細析出物を形成するが、０．１５％を超えると析出物が粗大化するようになるため、０．１５％以下とする。より好ましくは０．１０％以下である。
【００４８】
Ｗ
ＷはＴｉと微細析出物を形成するが、１．５％を超えると析出物が粗大化するようになるため、１．５％以下とする。より好ましくは１．０％以下である。
【００４９】
これらの元素の添加においては、Ｃ、Ｔｉ、Ｍｏ、Ｎｂ、Ｖ、Ｗの原子比を規定することが炭化物の微細化に有効で（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）＋（Ｖ／５１）＋（Ｗ／１９２）｝を０．５以上、１．５以下とした場合、粒径１０ｎｍ未満の微細析出物の形成が容易となる。さらに好ましくは、０．７以上、１．３以下である。
【００５０】
Ｎｂ、Ｖ、Ｗの一種または二種以上を含む微細な炭化物の場合は、（Ｔｉ＋Ｎｂ＋Ｖ）／（Ｍｏ＋Ｗ）が０．２〜２．０、更に微細な炭化物の場合は０．７〜１．５であることが観察された。
【００５１】
また、本発明鋼では、上記添加元素以外の残部はＦｅ及び不可避的不純物とするが、脱酸材としてＡｌを０．１％以下添加することができる。
【００５２】
強度、延性を向上させる場合、Ｎｉ、Ｃｒの一種または二種をＮｉ≦２％、Ｃｒ≦２％の範囲で添加してもかまわない。
【００５３】
熱間鍛造部品の靭性を向上させる場合、不可避不純物であるＰ≦０．０４０％、Ｎ≦８０ｐｐｍに規制することが望ましい。
【００５４】
尚、これらの元素の含有量や添加の有無により本発明の効果が損なわれることはない。
【００５５】
３．製造条件
図１は本発明に係る熱間鍛造部品の概略製造工程図でＳ１は棒鋼製造工程、Ｓ２は搬送工程、Ｓ３は製品仕上げ過程を示す。棒線材製造工程（Ｓ１）で鋼塊を熱間圧延し棒鋼とし、製品仕上げ過程（Ｓ３）で棒鋼を熱間鍛造加工し、所望の部品形状とした後、析出処理で微細析出物を析出させ引張強さ７００ＭＰａ以上とする。（本工程において、熱間鍛造後の冷却速度を調整し、析出処理を省略することも可能である。）
【００５６】
以下に製品仕上げ過程における望ましい製造工程について詳細に説明する。
【００５７】
鍛造加熱温度
鍛造加熱温度は１１００℃以上とする。本発明では、鍛造後の析出処理などにより微細析出物を析出させるため、鍛造時に溶解時から残存する炭化物を固溶させる。
【００５８】
鍛造加熱温度を１１００℃未満とした場合、溶解時から残存するＴｉ−Ｍｏ系炭化物等が固溶しないため１１００℃以上とする。
【００５９】
鍛造温度
鍛造温度は１２００℃以下、９００℃以上とする。鍛造後の析出処理により微細析出物を析出させるが、その析出促進のために１２００℃以下の温度で鍛造する必要がある。また、９００℃未満になると鍛造荷重が高くなりすぎるために熱間では鍛造が難しい。よって、鍛造温度は９００℃以上１２００℃以下に限定する。
【００６０】
鍛造後の冷却速度の調整または、冷却後の析出処理により所望のミクロ組織とする。冷却速度の調整による場合は、微細析出物の析出温度範囲の７００〜５５０℃を、微細析出物が得られる限界冷却速度（０．５℃／ｓｅｃ）以下で冷却する。０．５℃／ｓｅｃ．以下で冷却するためには、炉に入れて炉冷することが望ましい。
【００６１】
また、炉の温度を５５０℃〜７００℃の温度に保持し、１０分間以上６０分以下で保持することも可能である。この場合、母相はフェライト単相組織となる。また、保持時間は６０分で充分フェライト単相となるため、６０分を超えて保持しても変わらない。よって、１０分以上６０分以下で保持する。
【００６２】
再加熱による析出処理では、母相をフェライト単相とし、強度向上に寄与する微細析出物を析出させることが必要で、加熱温度はベイナイトが生成しないよう５５０℃以上とし、７００℃を超えると析出物が粗大化するため５５０〜７００℃とする。
【００６３】
また、微細なＴｉ、Ｍｏなどの炭化物を生成、析出させるため当該温度域において１０分以上保持する。この場合、母相はフェライト単相組織となる。また、保持時間は６０分で充分フェライト単相となるため、６０分を超えて保持しても変わらない。よって、１０分以上６０分以下で保持する。
【００６４】
【実施例】
[実施例１]
表１に示す組成の鋼（Ｎｏ．１〜１４）について、熱間鍛造後の冷却速度の調整、再加熱処理を種々の条件で行い、強度、靭性および被削性について調査した。
【００６５】
供試鋼を１５０ｋｇ真空溶解炉にて溶製し、８０ｍｍ径の丸棒に鍛伸し、その後、種々の温度で加熱し、熱間鍛造により５０ｍｍ径とした。
【００６６】
Ｎｏ．１は熱間鍛造後、０．３℃／ｓｅｃ、Ｎｏ．８は熱間鍛造後、１５℃／ｓｅｃで室温まで冷却し、その他は熱間鍛造後、加熱炉で１５分保持し、室温まで空冷した。
【００６７】
その後、各供試材から、引張試験片、衝撃試験片を採取した。引張試験はＪＩＳ４号試験片により常温での強度特性を、衝撃試験はＪＩＳ３号のＵノッチ衝撃試験片により試験温度２０℃での吸収エネルギーを求めた。
【００６８】
組織観察は断面を光学顕微鏡で観察するとともに、析出物を透過型電子顕微鏡（ＴＥＭ）で観察し、その組成をエネルギー分散型Ｘ線分光装置（ＥＤＸ）により求めた。
【００６９】
表２に試験結果を示す。Ｎｏ．１〜７は本発明例でフェライト単相組織に１０ｎｍ以下の微細析出物を有し、引張強さ７００ＭＰａ以上、降伏比０．８５以上で１２０Ｊ／ｃｍ²以上の優れた靭性が得られている。
【００７０】
一方、Ｎｏ．８〜１３は比較例で、Ｎｏ．８は鍛造後の冷却速度が速く、微細析出物が析出せず、母相もベイナイト組織のため、強度、降伏比ともに本発明範囲外で靭性も低い。
【００７１】
Ｎｏ．９は鍛造後の保持温度が高く、析出物が粗大化し、母相中にパーライトが析出し、降伏比が本発明範囲外で、衝撃値も低い。
【００７２】
Ｎｏ．１０は鍛造後の保持温度が低いためベイナイトが析出し、微細析出物による析出強化量が低下し、降伏強度が低く、降伏比が本発明範囲外で、衝撃値も低い。
【００７３】
Ｎｏ．１１は鍛造加熱温度（鍛造前加熱温度）が低く、Ｔｉ、Ｍｏ系炭化物が固溶せず、そのため析出物粒径が大きく、降伏比が本発明範囲外で低い。
【００７４】
Ｎｏ．１２は高Ｃ系組成のため、母相組織がフェライト・パーライトで、析出物粒径、降伏比が本発明範囲外である。
【００７５】
Ｎｏ．１３はＭｏ無添加系の組成のため、母相組織がフェライト・マルテンサイトで、降伏比が本発明範囲外である。
【００７６】
Ｎｏ．１４は鍛造温度が１２００℃よりも高いために、充分な析出物が得られず、降伏比が本発明範囲外である。
【００７７】
【表１】

【００７８】
【表２】

【００７９】
[実施例２]
表３に示す組成の鋼（Ｎｏ．１５〜１８）について、熱間鍛造後の冷却速度の調整、再加熱処理を種々の条件で行い、強度、靭性および被削性について調査した。
【００８０】
供試鋼を１５０ｋｇ真空溶解炉にて溶製し、８０ｍｍの径の丸棒に鍛伸し、その後、Ｎｏ．１５〜１７は１２５０℃で加熱し、１１００℃にて熱間鍛造を行い５０ｍｍ径とした後、析出処理（６００℃−１０分保持、空冷）した。
【００８１】
Ｎｏ．１８（従来材）は、１２５０℃で加熱し、熱間鍛造により５０ｍｍ径とし、空冷後、焼入れ焼戻し処理を行った。
【００８２】
その後、各供試材について強度、靭性および被削性を調査した。引張試験はＪＩＳ４号試験片により常温での強度特性を、衝撃試験はＪＩＳ３号のＵノッチ衝撃試験片により試験温度２０℃での吸収エネルギーを求めた。
【００８３】
被削性をドリル切削試験（試験片：５０ｍｍφ、２０ｍｍ厚）により評価した。ＪＩＳ高速度工具鋼ＳＫＨ５１の６ｍｍφのストレートドリルで、送り０．１５ｍｍ／ｒｅｖ、回転数７４５ｒｐｍ、１断面当たり３０箇所の貫通穴を開け、ドリルが切削不能になるまでの総穴数で評価した。
【００８４】
表４に試験結果を示す。本発明例のＮｏ．１５は引張強さ、被削性ともに従来例（Ｎｏ．１８）と同等であり、非調質材であっても調質材に匹敵する優れた特性が得られている。
【００８５】
Ｎｏ．１６、１７は、降伏強度は従来材（Ｎｏ．１８）と同等であるが、鋼組成が本発明の範囲外のためミクロ組織が本発明を満足せず、引張強さが高く被削性に劣る。
【００８６】
【表３】

【００８７】
【表４】

【００８８】
【発明の効果】
本発明によれば、熱間鍛造後、調質処理を行うことなく調質処理材と同等の強度、被削性を有する引張強さ７００ＭＰａ以上の機械構造用非調質型熱間鍛造部品およびその製造方法が得られ、産業上極めて有用である。
【図面の簡単な説明】
【図１】本発明鋼の製造工程の一例を示す図である。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a non-tempered hot-forged part for machine structure and a method for producing the same, and in particular, even when non-tempered after hot forging, the tensile strength is 700 MPa or higher and the yield ratio is 0.85 or higher, It relates to a material having excellent fatigue characteristics and machinability.
[0002]
[Prior art]
As structural parts used in automobiles and construction machinery, there are tempered parts that are carbon steel for mechanical structures and alloy steels for mechanical structures, hot-forged into parts and then quenched and tempered, and non-tempered parts that omit quenching and tempering. It is used.
[0003]
In the case of non-tempered parts, because of the ferrite-pearlite two-phase structure, many tensile strengths are as low as 800 MPa or less, and strength improvement has been an issue.
[0004]
However, in increasing the strength of non-tempered parts, when yield strength is increased and fatigue strength is comparable to tempered parts, there is concern that the steel will harden and toughness and machinability will deteriorate, so the yield ratio In other words, it is necessary to increase the yield strength but not the tensile strength.
[0005]
Japanese Patent Application Laid-Open No. 11-29842 relates to a technique for improving the tensile strength and fatigue strength in a two-phase structure of ferrite and pearlite. However, since the C content is increased in the steel composition, there is a concern about a decrease in toughness.
[0006]
Japanese Laid-Open Patent Publication No. 7-109518 proposes that low-C steel has a dual phase structure of ferrite and bainite and is subjected to aging heat treatment after forging in order to improve the strength and yield ratio of non-heat treated steel. According to this technique, although the tensile strength is improved, the yield ratio is as low as 0.8 or less.
[0007]
Japanese Unexamined Patent Publication No. 2000-129393 proposes that the steel composition is a low carbon-high alloy system to improve the hardenability and a martensite bainite structure in order to improve the strength and toughness of the non-heat treated steel for hot forging. ing. However, a yield ratio lower than that of ferritic / pearlitic steel is described, and there is a concern about a decrease in fatigue strength.
[0008]
Japanese Patent Laid-Open No. 7-173531 proposes to perform tempering with a bainite-martensite structure after hot forging in order to improve the yield ratio of non-tempered steel. It doesn't reach.
[0009]
[Problems to be solved by the invention]
As described above, a steel material for non-tempered parts that satisfies all of toughness, yield strength and yield ratio comparable to tempered parts has not been developed yet.
[0010]
Therefore, an object of the present invention is to provide a machine structure non-tempered hot forged part for machine structure having a performance comparable to that of a tempered part in these characteristics, and a manufacturing method thereof.
[0011]
[Means for Solving the Problems]
The inventors of the present invention have made various investigations from the viewpoint of the composition of the components and production conditions of steel materials that are non-tempered materials and have strength and toughness comparable to those of tempered materials, and that do not impair machinability. When strengthening by precipitation strengthening using a material, the yield strength is increased but the yield ratio is high, so that the steel is not excessively hardened and the toughness and machinability are not impaired.
[0012]
The present invention has been made based on the above findings and further studies. That is, the present invention
1. Steel composition is mass%, C ≦ 0.15%, Si ≦ 1%, Mn ≦ 2%, Ti: 0.03-0.35%, Mo: 0.05-0.8%, balance Fe and It consists of inevitable impurities, has a ferrite single phase structure, and the proportion of fine precipitates with a particle size of less than 10 nm in the ferrite phase is 90% or more of the total precipitates, and the following formula (1)
0.5 ≦ (C / 12) / {(Ti / 48) + (Mo / 96)} ≦ 1.5— (1)
However, in said formula (1), let each element be content (mass%).
A non-tempered hot-forged part for machine structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.
[0013]
2. The fine precipitate is a carbide of Ti and Mo, wherein the tensile strength of the test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more. A non-tempered hot forged part for machine structure .
[0014]
3. The steel composition contains, in mass%, C ≦ 0.15%, Si ≦ 1%, Mn ≦ 2%, Ti: 0.03 to 0.35%, Mo: 0.05 to 0.8%, Furthermore, it contains one or two or more of Nb ≦ 0.08%, V ≦ 0.15%, W ≦ 1.5% by mass%, and the following formula (2)
0.5 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)} ≦ 1.5 --- (2)
However, in said formula (2), each element shall be content (mass%), and what is not contained shall be 0.
A non-tempered hot-forged part for machine structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.
[0015]
4). The tensile strength of a test piece cut out from an arbitrary position of a part according to claim 3, wherein the fine precipitate is a carbide containing Ti, Mo, and at least one of Nb, V, and W. Non-tempered hot forged parts for machine structures having a thickness of 700 MPa or more and a yield ratio of 0.85 or more.
[0020]
5. After the steel according to

claim

1 or 3 is heated to 1100 ° C. or higher, the steel composition is hot forged into a part shape at a temperature of 900 ° C. or higher and 1200 ° C. or lower. Non-tempered type for mechanical structure , characterized by cooling at a cooling rate of ℃ / sec or less, wherein a tensile strength of a test piece cut out from an arbitrary position of a part is 700 MPa or more and a yield ratio is 0.85 or more Manufacturing method for hot forged parts.
[0021]
6). The steel according to claim 1 or 3 having a steel composition heated to 1100 ° C or higher and then hot forged at a temperature of 900 ° C to 1200 ° C, and then cooled to 550-700 ° C for 10 minutes to 60 minutes. Manufacture of non-tempered hot forged parts for mechanical structures having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more of a test piece cut out from an arbitrary position of the part, characterized by being held below Method.
[0022]
7). After the steel according to

claim

1 or 3 is heated to 1100 ° C or higher, the steel composition is hot forged at a temperature of 900 ° C or higher and 1200 ° C or lower, heated to 550 to 700 ° C and held for 10 minutes or longer and 60 minutes or shorter. A method for producing a non-tempered hot forged part for mechanical structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.
[0023]
DETAILED DESCRIPTION OF THE INVENTION
The microstructure, component composition and production conditions of the present invention will be described in detail below.
[0024]
1. Microstructure The hot forged part according to the present invention has a microstructure in which fine precipitates having a particle size of less than 10 nm are dispersed and precipitated in a ferrite single-phase structure.
When a ferrite single-phase structure is used, toughness comparable to that of the tempered material is obtained, and when fine precipitates are dispersed and precipitated in the structure, strength properties comparable to the tempered material (high yield strength, high yield ratio) Is obtained. Fine precipitates are precipitated by adjusting the cooling rate after hot forging or after hot forging and by precipitation treatment.
[0025]
In the present invention, the fine precipitate has a particle size of less than 10 nm. When the particle size of the precipitate is 10 nm or more, the tensile strength of 700 MPa or more necessary for a machine structural component for transportation equipment and construction equipment cannot be obtained.
[0026]
In addition, when a fine precipitate having a particle size of less than 10 nm is deposited in a ferrite single phase structure, the yield ratio is 0.85 or more, and the increase in tensile strength with respect to the increase in yield strength, that is, the hardening of steel is small. Machinability comparable to tempered steel is obtained. If either the parent phase or the precipitate falls outside the scope of the present invention, the yield ratio is less than 0.85.
[0027]
The smaller the particle size of the fine precipitates, the more effective for improving the strength, preferably 5 nm, more preferably 3 nm or less, and carbides containing Ti and Mo as such fine precipitates, and further Nb, V, Carbides containing one or more of W are preferred.
[0028]
Although the distribution form of these fine precipitates is not particularly defined, it is desirable to uniformly disperse (disperse precipitation) in the matrix.
[0029]
In the present invention, if the size of fine precipitates is satisfied at 90% or more of the total precipitates, a target tensile strength of 700 MPa or more can be obtained. However, a precipitate having a size of 10 nm or more consumes a precipitate-forming element and adversely affects the strength.
[0030]
Although the effect of the present invention is not impaired even if a small amount of Fe carbide is contained in addition to the precipitate described above, toughness is inhibited when a large amount of Fe carbide having an average particle size of 1 μm or more is contained, in the present invention The upper limit of the size of Fe carbide contained is preferably 1 μm, and the content is preferably 1% or less of the whole.
[0031]
The ratio of the fine precipitates to the total precipitates is obtained by the following method. An electron microscope sample is prepared by an electropolishing method using a twin jet method and observed at an acceleration voltage of 200 kV. At that time, the crystal orientation of the parent phase is controlled so that the fine precipitates have a measurable contrast with respect to the parent phase, and the defocus is shifted from the normal focus in order to minimize the counting of the precipitates. Observe by method.
[0032]
Moreover, the thickness of the sample in the region where the precipitate particles are measured is evaluated by measuring the elastic scattering peak and the inelastic scattering peak intensity using electron energy loss spectroscopy.
[0033]
By this method, the measurement of the number of particles and the measurement of the sample thickness can be executed for the same region. The measurement of the number of particles and the particle diameter is carried out for four locations of a 0.5 × 0.5 μm region of the sample, and the precipitates distributed per 1 μm ² are calculated as the number for each particle diameter.
[0034]
From this value and the sample thickness, the number of precipitates distributed per 1 μm ^{3 of} particle diameter is calculated, and the ratio of the precipitates having a diameter of less than 10 nm to the measured total precipitates is calculated.
[0035]
Further, in the present invention, the ferrite single-phase structure is a ferrite area ratio of 95% or more, preferably 98% or more in cross-sectional structure observation (200-fold optical microscope structure observation).
[0036]
2. Component composition The steel of the present invention can achieve the desired performance with the microstructure described above, but the following component composition is preferred.
[0037]
C
If C is contained in an amount exceeding 0.15%, fine precipitates are coarsened and the strength is reduced. More preferably, it is 0.03% or more and 0.12% or less.
[0038]
Si
Si is added for deoxidation, but if it exceeds 1%, it dissolves in ferrite and the deformation resistance during cold working increases, so it is made 1% or less. More preferably, it is 0.15% or less.
[0039]
Mn
Mn is added because it is effective in improving the strength. However, if it exceeds 2%, the cold workability deteriorates, so the content is made 2% or less. More preferably, it is 0.5% or more and 1.8% or less.
[0040]
Ti
Ti is added for the purpose of precipitating fine precipitates containing Ti-Mo carbide together with Ti carbide and Mo and improving the strength. To ensure a tensile strength of 700 MPa or more, the content is set to 0.03% or more. On the other hand, when the content exceeds 0.35%, the precipitate is coarsened and the strength is decreased, so the content is set to 0.03 to 0.35%. More preferably, it is 0.03 to 0.20%.
[0041]
Mo
Mo is added for the purpose of precipitating fine precipitates containing Ti-Mo carbide together with Mo carbide and Ti and improving the strength. To ensure a tensile strength of 700 MPa or more, 0.05% or more. On the other hand, if added over 0.8%, a low-temperature transformation phase such as bainite is formed, precipitation strengthening due to fine precipitates is insufficient, and strength decreases. Therefore, the content is set to 0.05 to 0.8%. More preferably, it is 0.15 to 0.45%.
[0042]
Mo has a slow diffusion rate, and when it precipitates together with Ti, the growth rate of the precipitate is reduced, and a fine precipitate is easily obtained.
[0043]
(C / 12) / {(Ti / 48) + (Mo / 96)}
This parameter affects the size of the precipitate, and when it is 0.5 or more and 1.5 or less, formation of fine precipitates having a particle size of less than 10 nm is facilitated, which is preferable. More preferably, it is 0.7 or more and 1.3 or less.
[0044]
For fine Ti-Mo carbides, Ti and Mo in the carbides should be 0.2 ≦ Ti / Mo ≦ 2.0 in atomic ratio, and 0.7 ≦ Ti / Mo ≦ 1.5 for finer carbides. Was observed.
[0045]
Furthermore, when improving the characteristics, it is preferable to add one or more of Nb, V, and W.
[0046]
Nb
Nb forms fine precipitates together with Ti and contributes to an increase in strength. Further, the structure is refined and the ductility is improved by adjusting the crystal grains. If it exceeds 0.08%, it becomes too fine and the ductility is lowered, so it is made 0.08% or less. More preferably, it is 0.04% or less.
[0047]
V
V forms fine precipitates with Ti, but when it exceeds 0.15%, the precipitates become coarser, so 0.15% or less. More preferably, it is 0.10% or less.
[0048]
W
W forms fine precipitates with Ti, but if it exceeds 1.5%, the precipitates become coarse, so 1.5% or less. More preferably, it is 1.0% or less.
[0049]
In the addition of these elements, it is effective to define the atomic ratio of C, Ti, Mo, Nb, V, and W to refine the carbide (C / 12) / {(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 192)} is 0.5 or more and 1.5 or less, it becomes easy to form fine precipitates having a particle size of less than 10 nm. More preferably, it is 0.7 or more and 1.3 or less.
[0050]
In the case of a fine carbide containing one or more of Nb, V, and W, (Ti + Nb + V) / (Mo + W) is 0.2 to 2.0, and in the case of a finer carbide, 0.7 to 1.5. It was observed that
[0051]
In the steel of the present invention, the balance other than the additive elements is Fe and inevitable impurities, but 0.1% or less of Al can be added as a deoxidizer.
[0052]
When improving the strength and ductility, one or two of Ni and Cr may be added in the range of Ni ≦ 2% and Cr ≦ 2%.
[0053]
When improving the toughness of hot forged parts, it is desirable to regulate P ≦ 0.040% and N ≦ 80 ppm, which are inevitable impurities.
[0054]
In addition, the effect of this invention is not impaired by content of these elements, or the presence or absence of addition.
[0055]
3. Manufacturing Conditions FIG. 1 is a schematic manufacturing process diagram of a hot forged part according to the present invention. S1 is a steel bar manufacturing process, S2 is a conveying process, and S3 is a product finishing process. The steel ingot is hot-rolled into a steel bar in the bar wire manufacturing process (S1), and the steel bar is hot-forged in the product finishing process (S3) to obtain a desired part shape. The tensile strength is 700 MPa or more. (In this step, it is possible to adjust the cooling rate after hot forging and omit the precipitation treatment.)
[0056]
Hereinafter, a desirable manufacturing process in the product finishing process will be described in detail.
[0057]
Forging heating temperature The forging heating temperature is 1100 ° C. or higher. In the present invention, since fine precipitates are deposited by, for example, precipitation after forging, carbides remaining from the time of dissolution are solid-dissolved during forging.
[0058]
When the forging heating temperature is less than 1100 ° C., the Ti—Mo-based carbide remaining from the time of melting does not dissolve, so the temperature is set to 1100 ° C. or higher.
[0059]
Forging temperature The forging temperature is 1200 ° C. or lower and 900 ° C. or higher. Although fine precipitates are deposited by the precipitation treatment after forging, it is necessary to forge at a temperature of 1200 ° C. or less in order to promote the precipitation. Moreover, since it becomes too high forging load when it becomes less than 900 degreeC, forging is difficult in hot. Therefore, the forging temperature is limited to 900 ° C. or higher and 1200 ° C. or lower.
[0060]
A desired microstructure is obtained by adjusting the cooling rate after forging or by precipitation treatment after cooling. In the case of adjusting the cooling rate, the precipitation temperature range of 700 to 550 ° C. of the fine precipitate is cooled at a critical cooling rate (0.5 ° C./sec) or less at which the fine precipitate is obtained. 0.5 ° C./sec. In order to cool below, it is desirable to cool in a furnace.
[0061]
Moreover, it is also possible to hold | maintain the temperature of a furnace to the temperature of 550 degreeC-700 degreeC, and to hold | maintain for 10 minutes or more and 60 minutes or less. In this case, the parent phase has a ferrite single phase structure. Further, since the retention time is 60 minutes and the ferrite single phase is sufficient, even if the retention time exceeds 60 minutes, there is no change. Therefore, hold for 10 minutes or more and 60 minutes or less.
[0062]
In the precipitation treatment by reheating, it is necessary to make the parent phase a ferrite single phase and to precipitate fine precipitates that contribute to strength improvement. The heating temperature is set to 550 ° C. or higher so that bainite is not formed. Since the product becomes coarse, the temperature is set to 550 to 700 ° C.
[0063]
Moreover, in order to produce | generate and precipitate fine carbide | carbonized_materials, such as Ti and Mo, it hold | maintains for 10 minutes or more in the said temperature range. In this case, the parent phase has a ferrite single phase structure. Further, since the retention time is 60 minutes and the ferrite single phase is sufficient, even if the retention time exceeds 60 minutes, there is no change. Therefore, hold for 10 minutes or more and 60 minutes or less.
[0064]
【Example】
[Example 1]
About the steel (No. 1-14) of a composition shown in Table 1, adjustment of the cooling rate after hot forging and reheating process were performed on various conditions, and strength, toughness, and machinability were investigated.
[0065]
The test steel was melted in a 150 kg vacuum melting furnace, forged into an 80 mm diameter round bar, then heated at various temperatures, and hot forged to a 50 mm diameter.
[0066]
No. No. 1 after hot forging, 0.3 ° C./sec. No. 8 was cooled to room temperature at 15 ° C./sec after hot forging, and the others were held in a heating furnace for 15 minutes after hot forging and air cooled to room temperature.
[0067]
Thereafter, tensile test pieces and impact test pieces were collected from each test material. For the tensile test, the strength characteristics at room temperature were obtained with a JIS No. 4 test piece, and for the impact test, absorbed energy at a test temperature of 20 ° C. was obtained with a U-notch impact test piece of JIS No. 3.
[0068]
In the structure observation, the cross section was observed with an optical microscope, the precipitate was observed with a transmission electron microscope (TEM), and the composition was determined with an energy dispersive X-ray spectrometer (EDX).
[0069]
Table 2 shows the test results. No. 1 to 7 are examples of the present invention having fine precipitates of 10 nm or less in a ferrite single phase structure, and excellent toughness of 120 J / cm ² or more is obtained at a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more. .
[0070]
On the other hand, no. Nos. 8 to 13 are comparative examples. No. 8 has a fast cooling rate after forging, fine precipitates do not precipitate, and the parent phase is a bainite structure, so both strength and yield ratio are outside the scope of the present invention and low toughness.
[0071]
No. No. 9 has a high holding temperature after forging, the precipitates become coarse, pearlite is precipitated in the matrix, the yield ratio is outside the range of the present invention, and the impact value is low.
[0072]
No. No. 10 has a low holding temperature after forging, so that bainite precipitates, the precipitation strengthening amount due to fine precipitates decreases, the yield strength is low, the yield ratio is outside the range of the present invention, and the impact value is also low.
[0073]
No. No. 11 has a low forging heating temperature (heating temperature before forging), Ti and Mo-based carbides do not dissolve, so the precipitate particle size is large, and the yield ratio is low outside the scope of the present invention.
[0074]
No. Since No. 12 has a high C composition, the matrix structure is ferrite pearlite, and the precipitate particle size and yield ratio are outside the scope of the present invention.
[0075]
No. Since No. 13 is a Mo-free composition, the matrix structure is ferrite martensite and the yield ratio is outside the scope of the present invention.
[0076]
No. In No. 14, since the forging temperature is higher than 1200 ° C., sufficient precipitates cannot be obtained, and the yield ratio is outside the range of the present invention.
[0077]
[Table 1]

[0078]
[Table 2]

[0079]
[Example 2]
The steel (Nos. 15 to 18) having the composition shown in Table 3 was subjected to adjustment of the cooling rate after hot forging and reheating treatment under various conditions, and the strength, toughness and machinability were investigated.
[0080]
The test steel was melted in a 150 kg vacuum melting furnace and forged into a 80 mm diameter round bar. Nos. 15 to 17 were heated at 1250 ° C., hot forged at 1100 ° C. to obtain a diameter of 50 mm, and then subjected to precipitation treatment (600 ° C. for 10 minutes, air cooling).
[0081]
No. 18 (conventional material) was heated at 1250 ° C., made into a 50 mm diameter by hot forging, subjected to quenching and tempering treatment after air cooling.
[0082]
Thereafter, the strength, toughness and machinability of each specimen were investigated. For the tensile test, the strength characteristics at room temperature were obtained with a JIS No. 4 test piece, and for the impact test, the absorbed energy at a test temperature of 20 ° C. was obtained with a U-notch impact test piece of JIS No. 3.
[0083]
Machinability was evaluated by a drill cutting test (test piece: 50 mmφ, 20 mm thickness). With a 6 mmφ straight drill of JIS high speed tool steel SKH51, feed was 0.15 mm / rev, rotation speed was 745 rpm, 30 through holes were made per cross section, and the total number of holes until the drill became uncut was evaluated.
[0084]
Table 4 shows the test results. No. of the example of the present invention. No. 15 is equivalent to the conventional example (No. 18) in both tensile strength and machinability, and excellent characteristics comparable to the tempered material are obtained even with the non-tempered material.
[0085]
No. Nos. 16 and 17 have the same yield strength as the conventional material (No. 18), but the steel composition is outside the scope of the present invention, so the microstructure does not satisfy the present invention, and the tensile strength is high and machinability is achieved. Inferior.
[0086]
[Table 3]

[0087]
[Table 4]

[0088]
【The invention's effect】
According to the present invention, after hot forging, a non-tempered hot-forged part for machine structure having a tensile strength of 700 MPa or more having the same strength and machinability as a tempered material without tempering treatment and The production method is obtained and is extremely useful in industry.
[Brief description of the drawings]
FIG. 1 is a diagram showing an example of a production process of steel according to the present invention.

Claims

Steel composition is mass%,
C ≦ 0.15%,
Si ≦ 1%,
Mn ≦ 2%,
Ti: 0.03-0.35%,
Mo: 0.05-0.8%
It consists of the remaining Fe and inevitable impurities, has a ferrite single phase structure, and the proportion of fine precipitates with a particle size of less than 10 nm in the ferrite phase is 90% or more of the total precipitates, and the following formula (1)
0.5 ≦ (C / 12) / {(Ti / 48) + (Mo / 96)} ≦ 1.5— (1)
However, in said formula (1), each element shall be content (mass%).
A non-tempered hot-forged part for machine structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.

The fine precipitate is a carbide of Ti and Mo, wherein the tensile strength of the test piece cut out from an arbitrary position of the part is 700 MPa or more, and the yield ratio is 0.85 or more. A non-tempered hot forged part for a machine structure .

Steel composition is mass%,
C ≦ 0.15%,
Si ≦ 1%,
Mn ≦ 2%,
Ti: 0.03-0.35%,
Mo: 0.05-0.8%
In addition, in mass%,
Nb ≦ 0.08%,
V ≦ 0.15%,
W ≦ 1.5%
Containing one or more of the following formula (2)
0.5 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)} ≦ 1.5 --- (2)
However, in said formula (2), each element shall be content (mass%), and what is not contained shall be 0.
A non-tempered hot-forged part for machine structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.

The tensile strength of a test piece cut out from an arbitrary position of a part according to claim 3, wherein the fine precipitate is a carbide containing Ti, Mo, and at least one of Nb, V, and W. Non-tempered hot forged parts for mechanical structures having a thickness of 700 MPa or more and a yield ratio of 0.85 or more.

After the steel according to claim 1 or 3 is heated to 1100 ° C. or higher, the steel composition is hot forged into a part shape at a temperature of 900 ° C. or higher and 1200 ° C. or lower. ° C. / sec cooling below the cooling rate, wherein the tensile strength of the test piece cut out from an arbitrary position of the component is more than 700 MPa, the non-heat treated type for machine structural yield ratio is 0.85 or more Manufacturing method for hot forged parts.

The steel according to claim 1 or 3 having a steel composition heated to 1100 ° C or higher and then hot forged at a temperature of 900 ° C to 1200 ° C, and then cooled to 550-700 ° C for 10 minutes to 60 minutes. Manufacture of non-tempered hot forged parts for mechanical structures having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more of a test piece cut out from an arbitrary position of the part, characterized by being held below Method.

After the steel according to claim 1 or 3 is heated to 1100 ° C or higher, the steel composition is hot forged at a temperature of 900 ° C or higher and 1200 ° C or lower, heated to 550 to 700 ° C and held for 10 minutes or longer and 60 minutes or shorter. A method for producing a non-tempered hot forged part for mechanical structure , wherein the tensile strength of a test piece cut out from an arbitrary position of the part is 700 MPa or more and the yield ratio is 0.85 or more.