JP2004003010A - Soft-nitrided parts excellent in fatigue characteristics and manufacturing method thereof - Google Patents

Abstract

The present invention relates to a steel for nitrocarburizing having excellent strength and fatigue properties and a method for producing the same.
Kind Code: A1 A ferrite single phase structure in which fine precipitates having a particle size of less than 10 nm are dispersed and deposited, and in terms of mass%, C ≦ 0.15%, Si ≦ 0.3%, Mn ≦ 2%, Ti : 0.03 to 0.35%, Mo: 0.05 to 0.8%, one or more of Nb ≦ 0.08%, V ≦ 0.15%, W ≦ 1.5%, S : 0.03-0.1%, Pb ≦ 0.2%, Ca ≦ 0.005%, B ≦ 0.02%, 0.5 ≦ (C / 12) / ｛(Ti / 48) + (Mo /96)+(Nb/93)+(V/51)+(W/192)｝≦1.5 A nitrocarburized component consisting of the balance Fe and unavoidable impurities. After heating the steel having the above composition at 950 to 1250 ° C, hot forging is performed, and in the subsequent cooling, 550 to 700 ° C is cooled at a rate of more than 0.5 ° C / sec to form a desired shape. Hold at 700 ° C. for 10 minutes or more.
[Selection diagram] Fig. 1

Description

但し、各元素は含有量（質量％）とする。
【００１５】
４．微細析出物がＴｉ、Ｍｏの炭化物であることを特徴とする１乃至３のいずれか一つに記載の疲労特性に優れた軟窒化部品。
【００１６】
５．鋼組成として、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する２記載の疲労特性に優れた軟窒化部品。
【００１７】
６．鋼組成として更に式（２）を満足することを特徴とする５記載の強度、疲労特性に優れた軟窒化部品。
０．５≦（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）
＋（Ｖ／５１）＋（Ｗ／１８４）｝≦１．５　　　　（２）
但し、各元素は含有量（％）とし、含まれないものは０とする。
【００１８】
７．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５または６に記載の疲労特性に優れた軟窒化部品。
【００１９】
８．鋼組成として更に質量％で、Ｓ：０．０３〜０．１％、Ｐｂ≦０．２％、Ｃａ≦０．００５％、Ｂ≦０．０２％の一種または二種以上を含有することを特徴とする２、３、５、６のいずれか一つに記載の疲労特性に優れた軟窒化部品。
【００２０】
９．２、３、５、６、８のいずれか一つに記載の組成の鋼を９５０〜１２５０℃で加熱後、仕上げ圧延温度８００℃以上で圧延し、その後の冷却において、７００〜５５０℃を０．５℃／ｓｅｃ超えで冷却することを特徴とする軟窒化部品用棒鋼の製造方法。
【００２１】
１０．２、３、５、６、８のいずれか一つに記載の組成の鋼を９５０〜１２５０℃で加熱後、仕上げ圧延温度８００℃以上で圧延し、その後の冷却において、７００〜５５０℃を０．５℃／ｓｅｃ超えで冷却した棒鋼を、所望の形状に成形後、５５０〜７００℃にて１０分間以上保持する軟窒化処理を行うことを特徴とする疲労特性に優れた軟窒化部品の製造方法。
【００２２】
【発明の実施の形態】
本発明に係る軟窒化部品のミクロ組織、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。
【００２３】
１．ミクロ組織
本発明に係る軟窒化部品は、▲１▼その素材（圧延ままの棒鋼）段階において優れた冷間鍛造性と▲２▼冷間鍛造後においては被削性に優れ、更に▲３▼軟窒化処理後においては調質処理材の軟窒化処理材（従来鋼：例えばＳＣＭ４３５鋼）に比して優れた疲労特性とまた、同等の強度、靭性が得られるよう、そのミクロ組織をフェライト単相で且つ粒径１０ｎｍ未満の微細析出物を含む組織に規定する。
【００２４】
母相をフェライト単相組織とすることにより、素材段階での冷間鍛造性、および冷間鍛造後の切削性を向上させ、更に該組織中に微細析出物を分散析出させることにより軟窒化後、強度とともに疲労特性を向上させる。
【００２５】
本発明では微細析出物の粒径は１０ｎｍ未満とする。析出物の粒径が１０ｎｍ以上の場合、軟窒化後の析出強化が不充分で調質材を素材としたものと比して強度が向上せず、また、強度特性において降伏比が上昇せず、疲労特性の向上が得られない。
【００２６】
微細析出物の粒径は小さいほど有効で、望ましくは５ｎｍ、更に望ましくは３ｎｍ以下で、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。
【００２７】
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。
【００２８】
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、焼戻し後目的とする引張強さが得られる。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。
【００２９】
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。
【００３０】
本発明における微細析出物の全析出物に占める割合は、次の方法で決定できる。まず電子顕微鏡試料を、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。
【００３１】
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、１μｍ^２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１μｍ^３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
本発明に係る軟窒化部品は上述したミクロ組織で目的とする性能が得られるが、以下の成分組成とすることが好ましい。
【００３７】
Ｃ
Ｃは強度確保のため添加する。０．１５％超えで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。より好ましくは０．０３％以上０．１２％以下である。
【００３８】
Ｓｉ
Ｓｉは脱酸のため添加するが、０．３％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０．３％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系炭化物や、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、軟窒化処理材の強度および疲労特性を向上させるため添加する。０．０３％未満では析出物量が少なく所望の強度及び疲労特性が得られないため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、そのような効果が低下するため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系炭化物や、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、軟窒化処理材の強度および疲労特性を向上させるため添加する。軟窒化処理材において所望の引張強度とし、疲労特性を向上させるため０．０５％以上とし、一方、０．８％を超えて添加するとベイナイト相を形成し、析出強化量が低下するため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。
【００４４】
微細なＴｉ−Ｍｏ系炭化物では、炭化物中のＴｉ、Ｍｏは原子比で２．０≧Ｔｉ／Ｍｏ≧０．２、更に微細な場合は１．５≧Ｔｉ／Ｍｏ≧０．７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と微細析出物を形成して強度向上に寄与する。また、組織を微細化し、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。
【００５０】
また、微細なＴｉ−Ｍｏ−（Ｎｂ、Ｖ、Ｗ）系炭化物では、炭化物中の各元素は原子比で２．０≧（Ｔｉ＋Ｎｂ＋Ｖ）／（Ｍｏ＋Ｗ）≧０．２、更に微細な炭化物では１．５≧（Ｔｉ＋Ｎｂ＋Ｖ）／（Ｍｏ＋Ｗ）≧０．７であることが観察された。
【００５１】
本発明鋼では、冷間鍛造後や軟窒化処理材の被削性を向上させる場合は、０．０３≦Ｓ≦０．１％とし、Ｐｂ≦０．２％、Ｃａ≦０．００５％、Ｂ≦０．０２％の一種以上を添加することができる。
【００５２】
また、本発明鋼では上記添加元素以外の残部はＦｅ及び不可避不純物とするが、脱酸剤としてＡｌを０．１％以下添加することができる。更に素材の冷間鍛造性を向上させる場合、Ｐ≦０．０４０％、Ｎ≦８０ｐｐｍとし、強度、延性を向上させる場合、Ｎｉ≦２％、Ｃｒ≦２％の一種以上を添加することができる。
【００５３】
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。
【００５４】
３．製造条件
図１は本発明に係る軟窒化部品の概略製造工程図でＳ１は素材となる棒鋼製造工程、Ｓ２は搬送工程、Ｓ３は製品（軟窒化部品）仕上げ工程を示す。棒鋼製造工程（Ｓ１）で鋼塊を熱間圧延し棒鋼とし品質検査後、出荷する。
【００５５】
製品（軟窒化部品）仕上げ工程（Ｓ３）で、該棒鋼を所定の寸法に切断し、冷間鍛造、冷間曲げなどの冷間鍛造を行い、必要に応じてドリル穿孔や旋削等の切削加工で所望の形状とした後、軟窒化処理を行い製品とする。
【００５６】
尚、冷間鍛造にかわり熱間鍛造を用いても良く、この場合は熱間鍛造後、冷間矯正が行われる場合がある。また、最終製品にペンキやメッキ等の皮膜処理がなされる場合もある。以下に望ましい製造工程について詳細に説明する。
【００５７】
圧延加熱温度
圧延加熱温度は９５０〜１２５０℃とする。本発明では、圧延材（素材となる棒鋼）に微細析出物が析出し冷間鍛造性を損なわないよう、熱間圧延時に溶解時から残存する炭化物を固溶させる。
【００５８】
圧延加熱温度は９５０℃未満とした場合、仕上までの圧延荷重が大きくなりすぎて圧延が不可能である。また、１２５０℃を超えると、析出物が再固溶するため、微細析出物が析出し、強度が上昇するため、冷間鍛造性が悪化する。よって、加熱温度は９５０℃〜１２５０℃とする。
【００５９】
圧延仕上げ温度
圧延仕上げ温度は８００℃未満では圧延荷重が高く真円度が劣化するため８００℃以上とする。
【００６０】
冷却速度
冷間鍛造前に微細析出物が析出し、冷間鍛造性を損なわないよう、圧延後の冷却速度を規定する。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）超えで冷却する。
【００６１】
軟窒化処理（析出処理）
得られた棒鋼を素材とし、冷間鍛造後、切削加工等により部品形状とする。その後、軟窒化処理を行う。軟窒化処理は微細析出物を析出させるように、加熱温度：５５０〜７００℃、保持時間１０分以上で行う。５５０℃未満では、十分な量の析出物が得られず、７００℃超えでは析出物が粗大化するため、５５０〜７００℃とする。
【００６２】
尚、冷間鍛造にかわり熱間鍛造を用いた場合、熱間鍛造後の冷間矯正や切削加工性の観点から、微細析出物が析出しないように熱間鍛造時の加熱温度を９５０〜１２５０℃、鍛造後の冷却速度を０．５℃／ｓｅｃ超えで行う。
【００６３】
【実施例】
表１に示す組成の鋼（Ｎｏ．１〜１１）を１５０ｋｇ真空溶解炉にて溶製し、圧延を１１００℃加熱、９２０℃仕上げで行い、その後１℃／ｓｅｃで室温まで冷却し３６ｍｍφの棒鋼とした。Ｎｏ．１１は従来材である。
【００６４】
Ｎｏ．１〜１０は圧延まま、Ｎｏ．１１は圧延後、球状化処理したものから冷間据えこみ加工用の試験片（直径１４ｍｍ、高さ／直径比：１．５）を採取し、同心円状溝突付きダイスを取りつけた拘束圧縮盤により圧縮加工時の変形抵抗、割れ発生限界加工率を平均歪速度０．０１／ｓｅｃで調査した。
【００６５】
変形抵抗は平均歪が１．５（圧下率７０％）の荷重を拘束係数と変形前の初期面積で除して求めた。割れ発生限界加工率は実際に割れの発生した圧下率とした。
【００６６】
尚、変形抵抗、割れ発生限界か効率は部品鍛造時の圧延工具寿命、割れ発生不良率と相関があることが知られている。
【００６７】
更に、被削性をドリル切削試験により評価した。上記圧縮試験採取材の残部を２５ｍｍ径まで冷間引抜き加工し、Ｎｏ．１〜１０は引抜き加工まま、Ｎｏ．１１は軟化焼鈍後、２０ｍｍ厚に切断したものを試験材として、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５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。
【００６８】
また、芯部の硬度（試験荷重１００ｇ）を圧延ままの素材（冷間鍛造前）、圧縮試験片（冷間鍛造後）について調査した。
【００６９】
軟窒化処理材はＮｏ．１〜１０は２５ｍｍ径まで冷間引抜き加工した鍛造材にガス軟窒化処理し、鋼Ｎｏ．１１は冷間引抜き加工後、焼入れ焼戻し処理を行ったものにガス軟窒化処理し製造した。ガス軟窒化処理はＮＨ_３：Ｎ_２：ＣＯ_２＝５０：４５：５の雰囲気で５２５〜７２５℃に加熱し、５時間保持して行った。
【００７０】
これらの軟窒化処理材について、組織観察し、芯部の硬度、引張特性および疲労特性を調査した。
【００７１】
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、組成をエネルギー分散型Ｘ線分光装置（ＥＤＸ）により求めた。
【００７２】
芯部の硬度は、冷間鍛造材と同じ位置にて測定した。また、ガス窒化後の表面硬さは表面から０．０２ｍｍの位置で測定を行い、有効硬化層深さはＨｖ４００となる表面からの深さと定義して測定した。
【００７３】
疲労特性は小野式回転曲げ疲労試験により疲労限を求めた。小野式回転曲げ疲労試験片（平行部８ｍｍφ）はφ２５ｍｍのものより採取し、上述した軟窒化処理を施した。
【００７４】
表２に試験結果を示す。Ｎｏ．１〜６が本発明に係る軟窒化部品、Ｎｏ．７〜１０が比較例、Ｎｏ．１１が既存の軟窒化用鋼による従来例である。
【００７５】
表から明らかなように、Ｎｏ．１〜６の軟窒化処理材はその素材（圧延材：棒鋼）の硬度、変形抵抗、限界加工率および冷間鍛造後の硬度、ドリル切削加工性も従来材の球状化処理材と同等である。更に、軟窒化処理後においては、従来材を調質処理後軟窒化処理したものと強度特性が同等で、疲労特性は更に優れている。
【００７６】
これに対しＮｏ．７〜１０は得られたミクロ組織が本発明範囲外で、強度、疲労特性に劣る。
【００７７】
Ｎｏ．７は軟窒化処理温度が本発明範囲外で高いため、軟窒化処理後後、フェライト＋パーライト組織となり、また、析出物の粒径も大きいため降伏強度が７００ＭＰａ以下と低い。
【００７８】
Ｎｏ．８は軟窒化処理温度が本発明範囲外で低いため、軟窒化処理後、ベイナイト組織となり、Ｃが固溶されたため、析出強化が不足し降伏強度が７００ＭＰａ以下と低い。
【００７９】
Ｎｏ．９はＣが高いため圧延材の硬度が高く、変形抵抗、限界加工率、ドリル寿命が従来材に及ばない。また、軟窒化処理後の析出物が大きく十分な析出強化が得られず、従来材より強度が低い。
【００８０】
Ｎｏ．１０はＳｉ、Ｍｎが高いため圧延材の硬度が高く、変形抵抗、限界加工率、ドリル寿命が従来材に及ばない。また、析出処理後、フェライト＋マルテンサイト組織となり、Ｃが固溶されたため微細析出物にもかかわらず十分な析出強化が得られず、従来材より強度が低い。またＮｏ．７〜１０のいずれの疲労限も従来材より低い。
【００８１】
【表１】

【００８２】
【表２】

【００８３】
【発明の効果】
本発明によれば、疲労特性に優れ且つ軟窒化処理後の冷間鍛造性、冷間鍛造後の被削性に優れた軟窒化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】本発明に係る軟窒化部品の製造工程の一例を示す図。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a nitrocarburized component and a method for producing the same, and more particularly to a nitrocarburized component having excellent strength and fatigue characteristics by adjusting a microstructure, which is preferable for automobiles and construction machines.
[0002]
[Prior art]
Car crankshafts and gears are required to have excellent fatigue properties and wear resistance. They are used after forming JIS standard material SCM435 etc. and then nitrocarburizing them. However, these steels have poor cold forgeability. Heat treatment such as soft annealing was required, and the machinability after cold forging was not sufficient.
[0003]
Further, in order to obtain a desired strength after the cold forging, it is necessary to perform quenching and tempering before soft nitriding, and it is inevitable that the production cost increases together with spheroidizing annealing for softening.
[0004]
Therefore, there is a strong demand for the development of a soft nitrided material having excellent cold forgeability, and various techniques have been proposed.
[0005]
Japanese Patent Application Laid-Open No. 9-279295 describes that the structure before soft nitriding is made into a ferrite bainite structure by adjusting the composition of the components to suppress a decrease in hardness after the treatment and further utilize precipitation strengthening of V. However, bainite has a high dislocation density and is inferior in cold forgeability before soft nitriding.
[0006]
Japanese Patent Application Laid-Open No. 2000-345259 discloses that the steel composition is made low in C and Mn to lower the hardness before soft nitriding and improve the cold forgeability. It is described that a compound is subjected to strain-induced precipitation to obtain a desired hardness after cold forging.
[0007]
However, since the amount of precipitation strengthening by the intermetallic compound is small and the hardness cannot be reduced before cold forging, it is inferior to cold forgeability and it is hard to say that the machinability after cold forging is sufficiently excellent. .
[0008]
[Problems to be solved by the invention]
As described above, the cold forgeability is improved without performing special heat treatment such as spheroidizing annealing, while having excellent machinability after cold forging, and furthermore, before soft nitriding, the desired strength and Steels that do not require heat treatment to do so have not yet been developed.
[0009]
Therefore, in the present invention, in order to achieve excellent cold forgeability and high strength after nitrocarburizing, a nitrocarburized component which can be produced at a low production cost without the conventionally required heat treatment, and a method for producing the nitrocarburized component The purpose is to provide.
[0010]
[Means for Solving the Problems]
The present inventors have conducted intensive studies on the effects of the structure and composition on the cold forgeability of steel, the machinability after cold forging and the strength after soft nitriding in order to solve the above problems, and A new manufacturing method was proposed in which the machinability after cold forging was improved by changing the steel structure to a ferrite single phase structure, and the strength after soft nitriding was improved by the precipitation of fine precipitates during soft nitriding.
[0011]
Furthermore, it has been found as a new finding that the yield ratio of finely precipitated steel is higher than that of tempered steel, and that excellent fatigue strength can be obtained.
[0012]
The present invention has been made by further study based on the above findings, that is, the present invention,
1. A nitrocarburized component having an excellent fatigue characteristic, wherein the component has a ferrite single phase structure and fine precipitates having a particle size of less than 10 nm are dispersed and precipitated in the ferrite phase.
[0013]
2. In mass%, C ≦ 0.15%, Si ≦ 0.3%, Mn ≦ 2%, Ti: 0.03-0.35%, Mo: 0.05-0.8%, balance Fe and inevitable A nitrocarburized component having excellent fatigue properties according to 1, which is made of impurities.
[0014]
3. 3. The nitrocarburized component having excellent fatigue properties according to 2, wherein the steel composition further satisfies the formula (1).

However, each element is the content (% by mass).
[0015]
4. 4. The nitrocarburized component having excellent fatigue characteristics according to any one of 1 to 3, wherein the fine precipitate is a carbide of Ti and Mo.
[0016]
5. 3. The nitrocarburized component having excellent fatigue properties according to 2, which further comprises one or more of Nb ≦ 0.08%, V ≦ 0.15%, and W ≦ 1.5% as steel composition by mass%.
[0017]
6. 6. The nitrocarburized component having excellent strength and fatigue properties according to 5, wherein the steel composition further satisfies the formula (2).
0.5 ≦ (C / 12) / ｛(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)｝ ≦ 1.5 (2)
However, each element is set to the content (%), and those not included are set to 0.
[0018]
7. 7. The nitrocarburized component having excellent fatigue properties according to 5 or 6, wherein the fine precipitate is a carbide containing Ti, Mo, Nb, V, and W.
[0019]
8. The steel composition further includes one or more of S: 0.03 to 0.1%, Pb ≤ 0.2%, Ca ≤ 0.005%, and B ≤ 0.02% by mass%. The nitrocarburized component having excellent fatigue properties according to any one of 2, 3, 5, and 6, characterized by the above-mentioned features.
[0020]
9.2 After heating the steel having the composition described in any one of 3, 5, 6, and 8 at 950 to 1250 ° C, rolling at a finish rolling temperature of 800 ° C or higher, and then cooling to 700 to 550 ° C. Is cooled at a rate of more than 0.5 ° C./sec.
[0021]
10.2,3,5,6,8 The steel of the composition described in any one is heated at 950-1250 degreeC, it rolls at 800 degreeC or more of finish rolling temperature, and 700-550 degreeC in the subsequent cooling. A steel bar cooled to 0.5 ° C./sec or more, formed into a desired shape, and then subjected to a nitrocarburizing treatment for maintaining the steel bar at 550 to 700 ° C. for 10 minutes or more. Manufacturing method.
[0022]
BEST MODE FOR CARRYING OUT THE INVENTION
The microstructure, component composition and manufacturing conditions of the nitrocarburized component according to the present invention will be described in detail below.
[0023]
1. Microstructure The nitrocarburized component according to the present invention has (1) excellent cold forgeability in its raw material (rolled steel bar) stage, (2) excellent machinability after cold forging, and (3) After the nitrocarburizing treatment, the microstructure is changed to ferrite single-strength so that excellent fatigue characteristics and equivalent strength and toughness can be obtained as compared with the nitrocarburized treated material (conventional steel: for example, SCM435 steel). It is defined as a structure containing a fine precipitate having a phase and a particle size of less than 10 nm.
[0024]
By forming the parent phase into a ferrite single-phase structure, the cold forgeability at the material stage, and the machinability after cold forging are improved, and further, fine precipitates are dispersed and precipitated in the structure. Improves fatigue properties as well as strength.
[0025]
In the present invention, the particle size of the fine precipitate is less than 10 nm. When the grain size of the precipitate is 10 nm or more, the precipitation strengthening after nitrocarburizing is insufficient, and the strength is not improved as compared with the case where the tempered material is used, and the yield ratio does not increase in the strength characteristics. And improvement in fatigue characteristics cannot be obtained.
[0026]
The smaller the particle size of the fine precipitate is, the more effective it is, preferably 5 nm, more preferably 3 nm or less. Such fine precipitates are carbides containing a complex of Ti and Mo, and also a kind of Nb, V and W. Alternatively, a carbide containing two or more kinds is preferable.
[0027]
Although the distribution form of these fine precipitates is not particularly defined, it is desirable that they are uniformly dispersed (dispersed and precipitated) in the parent phase.
[0028]
In the present invention, if the size of the fine precipitates satisfies 90% or more of all the precipitates, a desired tensile strength can be obtained after tempering. However, a precipitate having a size of 10 nm or more consumes a precipitate-forming element and adversely affects the strength.
[0029]
Although the effect of the present invention is not impaired even if a small amount of Fe carbide is contained separately from the above-mentioned precipitates, the toughness is impaired when a large amount of Fe carbide having an average particle size of 1 μm or more impairs toughness. It is desirable that the upper limit of the size of the Fe carbide contained is 1 μm, and the content is 1% or less of the whole.
[0030]
In the present invention, the ratio of the fine precipitates to the total precipitates can be determined by the following method. First, an electron microscope sample is prepared by an electrolytic polishing method using a twin jet method, and observed at an acceleration voltage of 200 kV.
[0031]
At this time, the crystal orientation of the parent phase is controlled so that the fine precipitate has a measurable contrast with respect to the parent phase, and the defocus is shifted from the normal focus to minimize the number of precipitates. Observe by method.
[0032]
In addition, the thickness of the sample in the region where the precipitation particles are measured is evaluated by measuring the elastic scattering peak and the inelastic scattering peak intensity using electron energy loss spectroscopy.
[0033]
According to this method, the measurement of the number of particles and the measurement of the sample thickness can be performed for the same region. The measurement of the number of particles and the particle diameter is performed on four portions of a 0.5 × 0.5 μm region of the sample, and the number of precipitates distributed per 1 μm ² is calculated as the number for each particle size.
[0034]
From this value and the sample thickness, the number of precipitates per particle size distributed per 1 μm ³ is calculated, and the ratio of the precipitates having a diameter of less than 10 nm to the measured total precipitates is calculated.
[0035]
In the present invention, the ferrite single-phase structure is defined as having a ferrite area ratio of 95% or more, and preferably 98% or more, when observed in a cross-sectional structure (observation with a 200-fold optical microscope).
[0036]
2. Component Composition Although the nitrocarburized component according to the present invention can achieve the desired performance with the above-mentioned microstructure, the following component composition is preferable.
[0037]
C
C is added to ensure strength. If the content exceeds 0.15%, the 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 0.3%, it is dissolved in ferrite and the deformation resistance during cold working increases, so it is made 0.3% or less. It is more preferably at most 0.15%.
[0039]
Mn
Mn is added because it is effective for improving the strength, but if it exceeds 2%, the cold workability deteriorates, so Mn is made 2% or less. More preferably, it is 0.5% or more and 1.8% or less.
[0040]
Ti
Ti is added to finely precipitate Ti-based carbides and precipitates containing Ti-Mo-based carbides together with Mo, and to improve the strength and fatigue properties of the nitrocarburized material. If it is less than 0.03%, the amount of precipitates is small and desired strength and fatigue properties cannot be obtained, so the content is made 0.03% or more. On the other hand, if it exceeds 0.35%, the precipitates become coarse, and such an effect is not obtained. The content is reduced to 0.03 to 0.35%. More preferably, it is 0.03 to 0.20%.
[0041]
Mo
Mo is added in order to finely precipitate Mo-based carbides and precipitates containing Ti-Mo-based carbides together with Ti, and to improve the strength and fatigue properties of the nitrocarburized material. The desired tensile strength of the nitrocarburized material is set to 0.05% or more in order to improve the fatigue properties. On the other hand, when added in excess of 0.8%, a bainite phase is formed and the amount of precipitation strengthening is reduced. 05 to 0.8%. More preferably, it is 0.15 to 0.45%.
[0042]
Mo has a low diffusion rate, and when precipitated 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 is preferably 0.5 or more and 1.5 or less, because it is easy to form a fine precipitate having a particle size of less than 10 nm, which is preferable.
[0044]
In fine Ti-Mo based carbides, the atomic ratio of Ti and Mo in the carbide is 2.0 ≧ Ti / Mo ≧ 0.2, and in the case of finer particles, 1.5 ≧ Ti / Mo ≧ 0.7. Was observed.
[0045]
In order to further improve the characteristics, it is preferable to add one or more of Nb, V and W.
[0046]
Nb
Nb forms fine precipitates with Ti and contributes to improvement in strength. Further, the structure is refined, and the ductility is improved by sizing the crystal grains. If it exceeds 0.08%, it is excessively fine and the ductility decreases, so the content is made 0.08% or less. More preferably, it is 0.04% or less.
[0047]
V
V forms fine precipitates with Ti, but if it exceeds 0.15%, the precipitates become coarse, so that V is set to 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 the content is made 1.5% or less. More preferably, it is 1.0% or less.
[0049]
In the addition of these elements, defining the atomic ratio of C, Ti, Mo, Nb, V, and W is effective for miniaturization of carbides, and is (C / 12) / ｛(Ti / 48) + (Mo / When 96) + (Nb / 93) + (V / 51) + (W / 192)} is 0.5 or more and 1.5 or less, it becomes easy to form a fine precipitate having a particle size of less than 10 nm.
[0050]
In a fine Ti-Mo- (Nb, V, W) -based carbide, each element in the carbide has an atomic ratio of 2.0 ≧ (Ti + Nb + V) / (Mo + W) ≧ 0.2, and 1 in finer carbide. It was observed that 0.5 ≧ (Ti + Nb + V) / (Mo + W) ≧ 0.7.
[0051]
In the steel of the present invention, when improving the machinability of the material after cold forging or nitrocarburizing treatment, 0.03 ≦ S ≦ 0.1%, Pb ≦ 0.2%, Ca ≦ 0.005%, One or more of B ≦ 0.02% can be added.
[0052]
In the steel of the present invention, the remainder other than the above-mentioned additional elements is Fe and inevitable impurities, but 0.1% or less of Al can be added as a deoxidizing agent. Further, when improving the cold forgeability of the material, P ≦ 0.040% and N ≦ 80 ppm, and when improving the strength and ductility, one or more of Ni ≦ 2% and Cr ≦ 2% can be added. .
[0053]
The effects of the present invention are not impaired by the content of these elements or the presence or absence of these elements.
[0054]
3. 1. Manufacturing Conditions FIG. 1 is a schematic manufacturing process diagram of a nitrocarburized component according to the present invention, wherein S1 is a bar steel manufacturing process as a raw material, S2 is a transport process, and S3 is a product (nitrocarburized component) finishing process. In the steel bar manufacturing process (S1), the steel ingot is hot-rolled into a steel bar and quality inspected before shipment.
[0055]
In the product (nitrocarburized component) finishing step (S3), the steel bar is cut into a predetermined size, and cold forging such as cold forging or cold bending is performed. If necessary, machining such as drilling or turning is performed. , And then nitrocarburizing to obtain a product.
[0056]
Note that hot forging may be used instead of cold forging, and in this case, cold straightening may be performed after hot forging. Further, the final product may be subjected to a coating treatment such as paint or plating. Hereinafter, a desirable manufacturing process will be described in detail.
[0057]
Rolling heating temperature The rolling heating temperature is 950 to 1250 ° C. In the present invention, carbide remaining from the time of melting at the time of hot rolling is solid-dissolved so that fine precipitates do not precipitate on the rolled material (steel bar as a raw material) and impair the cold forgeability.
[0058]
If the rolling heating temperature is less than 950 ° C., the rolling load to the finish becomes too large to perform rolling. On the other hand, when the temperature exceeds 1250 ° C., the precipitates form a solid solution again, so that fine precipitates precipitate and the strength increases, so that the cold forgeability deteriorates. Therefore, the heating temperature is 950 ° C to 1250 ° C.
[0059]
Rolling finishing temperature If the rolling finishing temperature is lower than 800 ° C, the rolling load is high and the roundness is deteriorated, so that the rolling finishing temperature is set to 800 ° C or higher.
[0060]
Cooling rate The cooling rate after rolling is specified so that fine precipitates are deposited before cold forging and the cold forgeability is not impaired. The fine precipitate is cooled at a deposition temperature range of 700 to 550 ° C. at a rate exceeding the limit cooling rate (0.5 ° C./sec) at which the fine precipitate is obtained.
[0061]
Soft nitriding treatment (precipitation treatment)
The obtained steel bar is used as a material, and after cold forging, it is shaped into a part by cutting or the like. Thereafter, a soft nitriding treatment is performed. The soft nitriding treatment is performed at a heating temperature of 550 to 700 ° C. and a holding time of 10 minutes or more so as to precipitate fine precipitates. If the temperature is lower than 550 ° C., a sufficient amount of precipitate cannot be obtained, and if the temperature exceeds 700 ° C., the precipitate is coarsened.
[0062]
When hot forging is used instead of cold forging, the heating temperature during hot forging is set to 950 to 1250 so that fine precipitates are not deposited from the viewpoint of cold straightening and cutting workability after hot forging. The cooling rate after forging is 0.5 ° C./sec.
[0063]
【Example】
A steel (Nos. 1 to 11) having the composition shown in Table 1 was smelted in a 150 kg vacuum melting furnace, rolled at 1100 ° C. and finished at 920 ° C., and then cooled to room temperature at 1 ° C./sec to obtain a 36 mmφ steel bar. And No. 11 is a conventional material.
[0064]
No. Nos. 1 to 10 were rolled. Reference numeral 11 denotes a constrained compression plate on which a test piece (diameter: 14 mm, height / diameter ratio: 1.5) for cold upsetting was sampled from the spheroidized product after rolling, and a concentric grooved die was attached. The deformation resistance at the time of compression working and the limit working rate of crack occurrence were investigated at an average strain rate of 0.01 / sec.
[0065]
The deformation resistance was determined by dividing a load having an average strain of 1.5 (a reduction ratio of 70%) by the constraint coefficient and the initial area before deformation. The crack generation limit working rate was defined as the rolling reduction at which cracks actually occurred.
[0066]
It is known that the deformation resistance and the crack occurrence limit or efficiency are correlated with the rolling tool life and the crack occurrence defect rate at the time of forging parts.
[0067]
Further, the machinability was evaluated by a drill cutting test. The remaining part of the compression test sample was cold drawn to a diameter of 25 mm. Nos. 1 to 10 were drawn as they were, No. 11 is a test material cut into a thickness of 20 mm after soft annealing and fed with a 6 mm φ straight drill of JIS high-speed tool steel SKH51 at a feed rate of 0.15 mm / rev, a rotation speed of 745 rpm, and five through holes per cross section. And the total number of holes until the drill could not be cut.
[0068]
In addition, the hardness of the core (test load: 100 g) was examined for the as-rolled material (before cold forging) and the compression test piece (after cold forging).
[0069]
The nitrocarburized material is No. Steel Nos. 1 to 10 were subjected to gas nitrocarburizing on a forged material cold-drawn to a diameter of 25 mm. Sample No. 11 was produced by subjecting a material subjected to quenching and tempering after cold drawing to gas soft nitriding. The gas soft nitriding treatment was performed by heating to 525 to 725 ° C. in an atmosphere of NH ₃ : N ₂ : CO ₂ = 50: 45: 5 and holding for 5 hours.
[0070]
The structure of the soft-nitrided material was observed, and the hardness, tensile properties and fatigue properties of the core were investigated.
[0071]
The structure was observed by observing the cross section with an optical microscope, observing the precipitate with a transmission electron microscope (TEM), and determining the composition with an energy dispersive X-ray spectrometer (EDX).
[0072]
The hardness of the core was measured at the same position as the cold forged material. The surface hardness after gas nitriding was measured at a position of 0.02 mm from the surface, and the effective hardened layer depth was defined as the depth from the surface at which Hv400 was obtained.
[0073]
Fatigue properties were determined by the Ono-type rotating bending fatigue test. Ono-type rotating bending fatigue test pieces (parallel portion 8 mmφ) were sampled from φ25 mm and subjected to the above-mentioned soft nitriding treatment.
[0074]
Table 2 shows the test results. No. Nos. 1 to 6 are nitrocarburized parts according to the present invention. 7 to 10 are comparative examples; Reference numeral 11 denotes a conventional example using the existing steel for nitrocarburizing.
[0075]
As is clear from the table, No. The nitrocarburized materials 1 to 6 have the same hardness, deformation resistance, critical working ratio, hardness after cold forging, and drilling workability of the material (rolled material: bar steel) as the conventional spheroidized material. . Further, after the nitrocarburizing treatment, the strength characteristics are the same as those obtained by subjecting the conventional material to the nitrocarburizing treatment after the tempering treatment, and the fatigue characteristics are further excellent.
[0076]
On the other hand, no. In Nos. 7 to 10, the obtained microstructure is out of the range of the present invention, and the strength and fatigue properties are poor.
[0077]
No. In No. 7, since the soft nitriding temperature is high outside the range of the present invention, a ferrite + pearlite structure is formed after the soft nitriding, and the yield strength is as low as 700 MPa or less due to the large grain size of the precipitate.
[0078]
No. Sample No. 8 had a low nitrocarburizing temperature outside the range of the present invention, and thus had a bainite structure after the nitrocarburizing process, and had a solid solution of C. Therefore, precipitation strengthening was insufficient and the yield strength was as low as 700 MPa or less.
[0079]
No. No. 9 has a high C, so that the hardness of the rolled material is high, and the deformation resistance, the critical working rate, and the drill life are inferior to those of the conventional material. Further, the precipitates after the nitrocarburizing treatment are so large that sufficient precipitation strengthening cannot be obtained and the strength is lower than that of the conventional material.
[0080]
No. Sample No. 10 has high hardness of rolled material due to high Si and Mn, and has less deformation resistance, critical working ratio and drill life than conventional materials. Further, after the precipitation treatment, a ferrite + martensite structure is formed, and since C is dissolved, sufficient precipitation strengthening cannot be obtained despite fine precipitates, and the strength is lower than that of the conventional material. No. Each of the fatigue limits of 7 to 10 is lower than the conventional material.
[0081]
[Table 1]

[0082]
[Table 2]

[0083]
【The invention's effect】
ADVANTAGE OF THE INVENTION According to this invention, the nitrocarburized component excellent in the fatigue characteristic and the cold forgeability after nitrocarburizing processing and the machinability after cold forging and its manufacturing method are obtained, and it is very useful industrially.
[Brief description of the drawings]
FIG. 1 is a view showing an example of a manufacturing process of a nitrocarburized component according to the present invention.

Claims (10)

A nitrocarburized component having an excellent fatigue characteristic, wherein the component has a ferrite single phase structure and fine precipitates having a particle size of less than 10 nm are dispersed and precipitated in the ferrite phase. In mass%, C ≦ 0.15%, Si ≦ 0.3, Mn ≦ 2%, Ti: 0.03-0.35%, Mo: 0.05-0.8%, balance Fe and unavoidable impurities The nitrocarburized component having excellent fatigue characteristics according to claim 1. The nitrocarburized component having excellent fatigue characteristics according to claim 2, wherein the steel composition further satisfies the formula (1).

However, each element is the content (% by mass). The nitrocarburized component having excellent fatigue characteristics according to any one of claims 1 to 3, wherein the fine precipitate is a carbide of Ti and Mo. 3. The nitrocarburized steel having excellent fatigue properties according to claim 2, further comprising one or more of Nb ≦ 0.08%, V ≦ 0.15%, and W ≦ 1.5% as a steel composition by mass%. parts. The nitrocarburized component having excellent strength and fatigue characteristics according to claim 5, wherein the steel composition further satisfies the formula (2).
0.5 ≦ (C / 12) / ｛(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)｝ ≦ 1.5 (2)
However, each element is set to the content (%), and those not included are set to 0. The nitrocarburized component having excellent fatigue characteristics according to claim 5 or 6, wherein the fine precipitate is a carbide containing Ti, Mo, Nb, V, and W. The steel composition further includes one or more of S: 0.03 to 0.1%, Pb ≤ 0.2%, Ca ≤ 0.005%, and B ≤ 0.02% by mass%. The nitrocarburized component having excellent fatigue characteristics according to any one of claims 2, 3, 5, and 6. A steel having a composition according to any one of claims 2, 3, 5, 6, and 8 is heated at 950 to 1250C, and then rolled at a finish rolling temperature of 800C or higher, and then cooled to 700 to 550C. Is cooled at a rate of more than 0.5 ° C./sec. The steel having the composition according to any one of claims 2, 3, 5, 6, and 8 is heated at 950 to 1250 ° C or higher, then rolled at a finish rolling temperature of 800 ° C or higher, and then cooled to 700 to 550. Nitrocarburizing excellent in fatigue characteristics, characterized in that a bar steel cooled at a temperature of more than 0.5 ° C / sec is formed into a desired shape and then subjected to a nitrocarburizing treatment for holding the steel bar at 550 to 700 ° C for 10 minutes or more. The method of manufacturing the part.

Images