JP4050512B2 - Manufacturing method of carburizing and quenching member and carburizing and quenching member - Google Patents
Manufacturing method of carburizing and quenching member and carburizing and quenching member Download PDFInfo
- Publication number
- JP4050512B2 JP4050512B2 JP2001391133A JP2001391133A JP4050512B2 JP 4050512 B2 JP4050512 B2 JP 4050512B2 JP 2001391133 A JP2001391133 A JP 2001391133A JP 2001391133 A JP2001391133 A JP 2001391133A JP 4050512 B2 JP4050512 B2 JP 4050512B2
- Authority
- JP
- Japan
- Prior art keywords
- mass
- carburized
- quenching
- carburizing
- quenched
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Expired - Fee Related
Links
Images
Landscapes
- Solid-Phase Diffusion Into Metallic Material Surfaces (AREA)
- Heat Treatment Of Articles (AREA)
Description
【０００１】
【発明の属する技術分野】
本発明は、疲労強度及び形状精度に優れた浸炭焼入れ部材の製造方法、ならびに浸炭焼入れ部材に関する。
【０００２】
【従来の技術】
浸炭焼入れは、鋼の表層の炭素濃度を高めた上で焼入れを行なう熱処理である。浸炭焼入れされた鋼部品は表層硬さが高く、かつ内部の靭性が高いので、疲労強度や靭性を要求される動力伝達機構部品などに適用されている。また、浸炭焼入れ前の素材の炭素量が比較的少ないため、切削や研磨といった機械除去加工や、圧延あるいは鍛造といった塑性加工が比較的容易であり、部品製造性にも優れているので、大量生産される自動車などの部品に多用されている。従って、浸炭焼入れ用鋼は、強度とともに部品の製造コストの低いことが前提として要求される。
【０００３】
浸炭焼入れされた部材の表層は、鋼としては最も高い硬度を示すため、表面近傍に微細な欠陥や脆弱部が存在すると、これらが疲労破壊や衝撃破壊の起点として作用するため、部材の強度が低下することにつながる。また、一般的な浸炭処理は、酸化性の変成ガス雰囲気中にて高温で行われるため、被処理部材の表面が酸化されやすい。このような表面酸化は、合金元素の枯渇による表層部の焼入性の低下を引き起こし、部材強度を劣化させることが知られている。これを改善する方法として、真空浸炭処理や、酸化を助長する元素を低減する一方非酸化性元素を添加した高強度鋼を適用することが行われている。また、耐摩耗性や耐ピッティング性が要求される場合には、高温強度や軟化抵抗を向上する元素の添加が有効であることが、特公平７−１１６５５１などに記載されている。
【０００４】
【発明が解決しようとする課題】
上述した浸炭焼入れ材に要求される強度とコストの両立は必ずしも十分ではない。特に、強度における疲労強度や衝撃強度と耐摩耗性や耐ピッティング性の両立、さらには、これらの強度と機械加工性や形状精度などの製造コストの両立は浸炭焼入れ部材における最も重要な課題であるが、従来の技術は今日の軽量化及び低廉化を追求する市場要求に十分応えていない。以下にその理由を説明する。
【０００５】
浸炭焼入れ部材の製造コストにおいて多くを占めるのは機械加工である。浸炭焼入れ前の切削加工性に対しては、素材硬度の影響が最も大きい。例えば前記した特公平７−１１６５５１などに記載されている高強度鋼は、一般的には高合金鋼となるため機械加工性に劣る。
【０００６】
例えば、疲労強度や衝撃強度（以下、これらを総称する場合は「破断強度」という）と、耐摩耗性や耐ピッティング性（以下、これらを総称する場合は「面疲労強度」という）に関しては、それぞれ単独であれば、鋼材化学成分の調整により向上させることが可能である。たとえば、破断強度に対してはＳｉ、Ｍｎ及びＣｒの低減とＮｉ及びＭｏの添加が有効であり、面疲労強度に対してはＳｉ、Ｃｒ及びＭｏの添加が有効であることが知られている。しかしながら、このような高合金化は、鋼材コストの上昇に加え、素材硬さの上昇により機械加工のコストを増大させたり、浸炭焼入れ時の浸炭性の劣化や残留オーステナイト組織の増大を招くため、限界がある。
【０００７】
他方、浸炭焼入れ部材の強度に影響する化学成分のうち、特にＳｉは破断強度に対しては最も有害である半面、面疲労強度には最も有益な元素である。従って、歯車などで要求されているような、破断強度と面疲労強度の両立を鋼材の化学成分のみで達成するには、ＮｉやＭｏなどの、Ｓｉ以外の比較的高価な強化元素を大量に添加することが必要になり、前述のように高合金化の弊害が大きくなるため、強度と製造性の両立ができない。
【０００８】
次に、部材に要求される形状や寸法の公差が厳しい場合は、浸炭焼入れ時に生ずる熱処理歪が大きいと、浸炭焼入れ後の状態では必要な寸法精度を確保できず、研削仕上げが必要となることがある。特に歯車などでは、形状精度が騒音や損傷の原因となるため、専用加工機による歯面研削を行なう場合がある。このような研削加工は一般的に切削加工よりもコストが高く、時に、浸炭焼入れ部材の全体製造コストの５０％以上を占める場合もある。したがって、浸炭焼入れ時においては、発生する歪がなるべく小さくなることが求められ、従来、主として焼入れ技術による改善が図られているが、十分な改善には至っていない。
【０００９】
また、浸炭焼入れにおける冷却媒体としてはほとんどの場合、油やソルトが用いられているが、ガス冷却による焼入れ（以下、「ガス焼入れ」と称する）も行われることもある。ガス焼入れは熱処理歪が小さいため、寸法や形状の精度を確保して加工コストを削減する観点において好都合である。しかしながら、ガスは冷却能が低いため、焼入組織が十分に形成されず不完全焼入れとなる場合がある。このような不完全焼入状態では、十分な強度や機械的特性が得られないので、結果的にガス冷却は小物部品にしか適用できなかった。
【００１０】
例えば、ガス焼入れの冷却能を向上せしめるには、冷却室のガス圧力を高くすることが有効であるが、設備の大型化や処理コストの増大を招く。このような冷却能の不足に対しては、鋼材の焼入性を高めることが有効であるが、鋼の焼入性を高める元素として代表的なＳｉ、Ｍｎ、Ｃｒ、Ｎｉ、Ｍｏなどは浸炭焼入れ前の素材の硬さを高める元素でもあり、機械加工性や冷間加工性に有害であり、部品の製造性を劣化させる。また、ＪＩＳ（日本工業規格）には、焼入性の高い肌焼鋼としてＳＮＣＭ６１６やＳＮＣＭ８１５などが規定されているが、これらは、最も一般的な肌焼鋼であるＳＣｒ４２０やＳＣＭ４２０に比べると、合金含有量が多いため焼ならしや焼なまし状態の硬さが非常に高く、切削加工や冷間加工が困難である。すなわち、ガス焼入れに適合する焼入性と、大量生産に適した部品の製造性を両立することは、本質的に困難だったのである。
【００１１】
本発明の課題は、浸炭焼入前の加工が容易であって、部材寸法に関係なくガス冷却による低歪で高精度の浸炭焼入が可能であり、しかも浸炭焼入れ後の破断強度や面疲労強度に優れ、部材全体の製造コストを大幅に削減できる浸炭焼入部材の製造方法と、それにより得られる浸炭焼入れ部材とを提供することにある。
【００１２】
【課題を解決するための手段及び作用・効果】
上記課題を解決するために、本発明の浸炭焼入れ部材の製造方法は、副成分の含有量が、Ｃ：０．１２〜０．２２質量％、Ｓｉ：０．４〜１．５質量％、Ｍｎ：０．２５〜０．４５質量％、Ｎｉ：０．５〜１．５質量％、Ｃｒ：１．３〜２．３質量％、Ｂ：０．００１〜０．００３質量％、Ｔｉ：０．０２〜０．０６質量％、Ｎｂ：０．０２〜０．１２質量％、Ａｌ：０．００５〜０．０５質量％、Ｍｏ：０．０５質量％以下に設定され、残部が主成分をなすＦｅと不可避不純物に設定され、また、９００℃から室温まで一定速度にて冷却したとき、冷却速度が少なくとも０．１℃／秒以下の範囲においてはベイナイトが生成せず、かつ、冷却速度が少なくとも１２℃／秒以上の範囲においてはフェライトが生成しなくなるように、前記各副成分の非浸炭状態における含有量が調整されてなる鋼により部材を構成し、１〜３０ＨＰａの減圧雰囲気にて前記部材を浸炭処理することにより、表面炭素濃度が０．６〜１．５質量％となる浸炭層を該部材に形成し、その浸炭処理後に前記部材を、焼入急冷度Ｈが０．０１〜０．０８となるよう不活性ガス又は窒素ガスを用いたガス冷却により焼入れするとともに、
前記ガス焼入れを行なう際のガス冷却雰囲気は、前記部材を構成する鋼の熱伝導率をλ（単位：ｋｃａｌ／ｍｈ℃）、該ガス冷却雰囲気における前記部材の表面熱伝達係数をα（単位：ｋｃａｌ／ｍｈ２℃）として、Ｈ≡０．５×（α／λ）にて定義される焼入急冷度Ｈが０．０１〜０．０８となる雰囲気が使用されることを特徴とする。
【００１３】
上記本発明の方法によると、浸炭焼入前の加工が容易であって、部材寸法に関係なくガス冷却による低歪で高精度の浸炭焼入が可能であり、しかも浸炭焼入れ後の破断強度や面疲労強度に優れ、部材全体の製造コストを大幅に削減できる。以下に、その理由を詳しく説明する。
【００１４】
本発明者らは、浸炭焼入れ材の破断強度、面疲労強度、素材硬度及び熱処理歪のすべてを同時に改善することを目的として、鋼材の化学成分及び熱処理につき、それらの相互作用も含めて研究した結果、以下のようなことを見出した。まず、面疲労強度の向上にはＳｉの活用が不可欠であるが、過剰な添加は破断強度に極めて有害に作用する。他方、浸炭焼入部材においては、被処理部材の表面が酸化されると、合金元素の枯渇による表層部の焼入性の低下を引き起こし、同様に破断強度の劣化を招く。そこで、減圧浸炭雰囲気又は不活性ガスを主体とする浸炭雰囲気にて部材を浸炭処理することにより、表面酸化の影響が大幅に軽減され、Ｓｉ添加により多かれ少なかれ生ずる破断強度低下のマージンを確保することができる。
【００１５】
次に、熱処理歪の低減に有効なガス焼入れを、広範囲な形状ないし寸法の部材に適用するには、鋼材の焼入性を高めることが有効である。本発明においては、減圧浸炭雰囲気又は不活性ガスを主体とする浸炭雰囲気を採用することを考慮して、９００℃から室温まで一定速度にて冷却したとき、冷却速度が少なくとも０．１℃／秒以下の範囲においてはベイナイトが生成せず、かつ、冷却速度が少なくとも１２℃／秒以上の領域ではフェライトが生成しなくなる鋼を採用する。
【００１６】
冷却速度の増加に伴い、鋼材の組織は一般的には以下の順に変化することが知られている：
Ｆ＋Ｐ →（Ｆ＋Ｐ＋Ｂ →Ｆ＋Ｂ →Ｂ →Ｂ＋Ｍ） →Ｍ
ここで、Ｆはフェライト、Ｐはパーライト、Ｂはベイナイト、Ｍはマルテンサイトを表す。なお、合金組成によっては単独ベイナイトとなる条件が存在しないこともありえる。また、パーライトはフェライトとセメンタイトとの共析組織であり、パーライトが生成することは、フェライトが生成するということでもある。そして、Ｆ＋Ｐ、あるいはＦ＋Ｂ等とあるのは、パーライト以外の形態で単独形成されるフェライトが存在することを意味する。上記冷却速度の範囲は、鋼の連続冷却変態線図（Continuous Cooling Transformation diagram：ＣＣＴ線図）を種々の冷却速度により測定することにより特定できる。
【００１７】
浸炭層の焼入れ硬化の要部を担うのは前述の通りマルテンサイトである。マルテンサイト変態は、よく知られている通り、原子の大きな拡散を伴うことなく結晶格子が擬剪断変形的に連携運動して生ずる。鋼のマルテンサイト変態は体積膨張が大きいことから、周囲の残留オーステナイトを大きく歪ませる形で進行する。これが鋼の焼入れ硬化の一因ともなる。しかし、変態に伴う歪みが大きいということは、硬さは向上する反面、焼入れ後に部材に寸法等の狂いを生じすいことも意味する。
【００１８】
本発明の場合、部材表面に浸炭層を形成し、その浸炭層に優先的にマルテンサイトを生成させることにより、表層部の硬度を増し、耐摩耗性を向上させるようにする。本発明では冷却による歪の増大を抑制するため、水や油よりも冷却速度が小さいガス冷却を採用している。従って、ガス冷却でも浸炭層に十分焼きが入るように、鋼の組成を、冷却速度が少なくとも１２℃／秒以上（以下、これを上限冷却速度という）の領域ではフェライトが生成しなくなるように設定する。１２℃／秒以上に冷却速度を大きくしてもフェライトが生成するようでは、ガス冷却によって浸炭層に十分にマルテンサイトが形成されず、硬さが不足することにつながる。
【００１９】
しかし、焼入れ性が過度に良好となるのも、本発明においては却って不利に作用する。すなわち、浸炭の影響が及ばない内層部においてもマルテンサイトが過剰に生成するようであれば、部材全体としてのマルテンサイト生成量が大きくなって寸法精度の低下につながる。そこで、ガス焼入時に、浸炭層においては十分にマルテンサイトが生成するが、内層部では過度のマルテンサイトが形成されないように、組成を選定することが重要である。具体的には、冷却速度が少なくとも０．１℃／秒以下（以下、下限冷却速度という）の範囲においてはベイナイトが生成しないようにする。０．１℃／秒以下の冷却速度でもベイナイトが生成するようであれば、浸炭層の影響が及ばない内層部にまで深く焼きが入って歪みが増大し、本発明の目的を達成できなくなる。
【００２０】
一方、本発明のもう一つの目的は、浸炭焼入前の加工を容易にできるようにすることであり、そのためには、焼きなまし状態、すなわちオーステナイトからの冷却速度が放冷または空冷に相当する範囲（以下、焼き鈍し冷却速度範囲という）において、加工性を向上させるのに十分低い素材硬度が得られなければならない。そこで、これを想定して０．１℃／秒未満の冷却速度に設定したとき、ベイナイトが生成しないように、非浸炭状態の鋼の組成を選定するのである。そこで、上記のように０．１℃／秒未満の冷却速度でベイナイトが生成しないようにすれば、実際の焼き鈍し冷却速度範囲にてベイナイトの生成が十分抑制され、フェライト＋パーライトの多い加工性に富んだ組織を得ることができる。
【００２１】
なお、内層部のフェライトあるいはベイナイト生成の限界冷却速度を実験的に決定する方法としては、鋼を焼きならし後、図４に示す形状及び寸法を有した試験片に加工する。この試験片には、片端から長手方向中央部まで有底孔が形成され、孔底に熱電対を溶接により固定する。この状態で、試験片を予め定められた一定流量の冷却ガス中に置き、さらに高周波誘導加熱によりオーステナイト化温度である９００℃に昇温する。温度が９００℃にほぼ一定に保持されれば、熱電対にて試験片中央位置の温度をモニタしながら高周波誘導加熱のパワーを制御しつつ減少させ、一定の冷却速度にて冷却する。冷却中、試験片は一端を固定し、他端位置に取り付けた支持棒の位置を差動トランスを用いてモニタすることにより、試験片寸法の温度変化をリアルタイム測定し、その寸法変化に現れる変曲点からフェライトあるいはベイナイト生成の温度を読み取る。また、冷却終了後、熱電対に対応する位置にて試験片を軸直交面にて切断し、光学顕微鏡観察によりフェライトあるいはベイナイトの有無を目視確認する。この測定を、冷却速度を種々に変化させて行い、フェライト及びベイナイトの有無をそれぞれ確認することにより、限界冷却速度を決定する。なお、この測定法では、フェライト及びベイナイトの生成を試験片寸法変化の変曲点と光学顕微鏡組織の両方にて確認するが、前者においては周知の微分解析による変曲点の決定には一定の誤差があり、その誤差の範囲で変態に対応する変曲点が認められなければ、フェライトあるいはベイナイト生成なしと判定する。他方、光学顕微鏡組織による観察では、断面上に０．１ｍｍ四方の観察視野をランダムに抜き出す形で９視野設定し、各視野にてフェライトあるいはベイナイトが目視にて観察されなければフェライトあるいはベイナイト生成なしと判定する。いずれの場合も、面積率にて１〜２％程度のフェライトあるいはベイナイトが生成していても、測定限界の問題から実質的にフェライトあるいはベイナイト生成なしとそれぞれ判定するものとする。
【００２２】
上記のように、熱処理歪と強度との双方の観点からガス焼入れに適した焼入性を得るには、合金元素の増量が有効であるが、高合金化による弊害は最小限に留める必要がある。一方、低い素材硬度を得るには、オーステナイトからの冷却速度が放冷または空冷に相当する範囲において、ベイナイトの生成を抑制し、フェライト・パーライトの多い組織とすることが必要である。焼入れ後の高い硬さと素材の低い硬さの両立は一般的には困難とされていたが、本発明においては、特有の合金元素の組み合わせにより、好ましい範囲で両立させることが可能となった。以下に詳細に説明する。
【００２３】
まず、上記のような効果を生み出す化学成分として、最も有効な元素はホウ素（Ｂ）である。Ｂはオーステナイトに固溶状態で存在する場合、冷却速度が２℃／秒以上の冷却において、フェライトやベイナイトの生成を抑制し焼入性を高めるが、冷却速度が１℃／秒以下においては、Ｂを添加しない場合と同じ組織及び硬さが得られる。
【００２４】
他方、焼入性や焼ならし硬さの調整は、Ｓｉ、Ｍｎ、Ｃｒ、Ｎｉ及びＭｏ等の添加によりなされる。ここで、一般的な肌焼鋼等においては、これら含有量の影響については、挙動は従来より十分に解析されており、ある程度の予測も可能であった。しかし、本発明では、低い素材硬度を得るためにＢの添加を必須とし、鋼材中には固溶Ｂが必ず存在する。このような固溶Ｂを有する鋼に関しては、上記合金元素の影響すなわちＢとの相互作用について、従来十分に明らかにされておらず、当然にそれらの適正な組成範囲についても何ら提案はなかった。
【００２５】
本発明者らは、この点に関して詳細に検討した結果、Ｎｉ及びＳｉが、固溶Ｂと共存する場合に焼入性の向上に対する寄与が特に大きく、反対に素材硬度に対する影響は少ないということを見出した。また、逆に、Ｍｎ及びＭｏは、固溶Ｂと共存する場合に、素材硬度を上昇させる影響が特に大きいことも判明した。本発明にて採用する鋼材組成は、Ｂ、Ｎｉ、Ｓｉの添加あるいは増量、及びＭｎ、Ｍｏの低減を基本として合金元素含有量を調整することにより、高い焼入性と低い素材硬度の両立を可能とした、特有のものである。
【００２６】
以下、本発明にて採用する鋼の組成限定理由について説明する。
（１）Ｃ：０．１２〜０．２２質量％
Ｃは浸炭焼入れ材の非浸炭部の強度を向上する元素である。しかし、含有量が０．１２質量％未満ではその効果が小さく、０．２２質量％を超えると素材硬度が高くなる。よって、Ｃの含有量は０．１２〜０．２２質量％とする。
【００２７】
（２）Ｓｉ：０．４０〜１．５０質量％
Ｓｉは面疲労強度の向上に有効な元素であるとともに、固溶Ｂと共存する場合に特に焼入性を向上し、素材硬度に対する影響が比較的小さい元素であるので、本発明において積極的に添加される元素である。その含有量が０．４０質量％未満では面疲労強度及び焼入性改善効果が小さく、１．５０質量％を超えるとＡ３変態点の上昇により、焼入れ前の均一オーステナイト化が困難になる。よって、Ｓｉの含有量は０．４０〜１．５０質量％とする。
【００２８】
（３）Ｍｎ：０．２５〜０．４５質量％
Ｍｎは焼入性を向上する元素であるが、０．２５質量％未満では効果が小さく、０．４５質量％を超えると固溶Ｂと共存する場合に素材硬度を上昇させる効果が特に大きく、部品製造性を著しく劣化させる。よって、Ｍｎの含有量は０．２５〜０．４５質量％とする。
【００２９】
（４）Ｎｉ：０．５０〜１．５０質量％
Ｎｉは浸炭鋼の強度や焼入性を向上する元素であり、固溶Ｂと共存すると特にその効果が大きい。しかし、含有量が０．５０質量％未満ではその効果が顕著でなく、１．５０質量％を超えるとベイナイト組織が生成しやすくなり、機械加工性を著しく劣化させる。よって、Ｎｉの含有量は０．５０〜１．５０質量％とする。
【００３０】
（５）Ｃｒ：１．３０〜２．３０質量％
Ｃｒは浸炭性を向上する元素であり、特にＳｉ及びＮｉによる浸炭性の劣化を防止する効果を有する元素である。しかし、含有量が１．３０質量％未満では効果が不足し、また、２．３０質量％を超えて添加すると、素材においてベイナイト組織が生成しやすくなり機械加工性を著しく劣化させる。よって、Ｃｒの含有量は１．３０〜２．３０質量％とする。
【００３１】
（６）Ｂ：０．００１０〜０．００３０質量％
Ｂは焼入性を著しく向上する元素であるが、０．００１０質量％未満では安定した効果が得られず、また、０．００３０質量％を超えて添加しても効果が飽和するので経済的でない。よって、Ｂの含有量は０．００１０〜０．００３０質量％とする。
【００３２】
（７）Ｔｉ：０．０２〜０．０６質量％
Ｔｉは窒化物を形成することにより、Ｂが窒化物となることを防止し、Ｂによる焼入性向上効果を安定させる元素である。しかし、その含有量が０．０２質量％未満では効果が小さく、０．０６質量％を超えて添加しても効果が飽和するので経済的でない。よって、Ｔｉの含有量は０．０２〜０．０６質量％とする。
【００３３】
（８）Ｎｂ：０．０２〜０．１２質量％
Ｎｂは結晶粒の成長を抑制し、強度を改善する上で有効な元素である。その含有量が０．０２質量％未満では効果が小さく、０．１２質量％を超えると凝固時に粗大な炭窒化物を形成して結晶粒成長抑制効果が減退する。よって、Ｎｂの含有量は０．０２〜０．１２質量％とする。
【００３４】
（９）Ａｌ：０．００５〜０．０５０質量％
Ａｌは溶製過程における脱酸を促進し酸化物系介在物量の低減に有効であるが、０．００５質量％未満では脱酸効果がなく、また、０．０５０質量％を超えて添加しても効果が飽和するので経済的でない。よって、Ａｌの含有量は０．００５〜０．０５０質量％とする。
【００３５】
なお、一般の浸炭用鋼に添加されることの多いＭｏは、Ｂと共存する場合に焼入性を向上する効果が特に大きい元素である。しかし、熱間加工上がりの状態や、焼ならし状態の素材硬度を著しく上昇させるので、製造性と強度の両立を目的とする本発明においては、積極的に添加する元素ではなく、好ましくは、採用する鋼のＭｏの含有量を０．０５質量％以下に制限するのがよい。
【００３６】
さらに、本発明において、浸炭層の表面炭素濃度は０．６〜１．５質量％とする。浸炭層の表面炭素濃度は浸炭焼入れ材の表面硬さに影響するが、０．６質量％未満では表面硬さが不足し、１．５質量％を超えると炭化物の析出量が多くなって、基地の焼入性が顕著に低下し、表面硬さが不足することにつながる。
【００３７】
また、浸炭雰囲気は窒素ガスが安価であり、本発明に好適に使用できるが、窒素ガスに代えてアルゴン等の不活性ガスを用いることもできる。
【００３８】
以下、本発明の浸炭焼入部材の製造方法において、さらに付加可能な要件について説明する。本発明者らにおいては１．５質量％までのＳｉを含有する鋼を採用するが、シリコン添加に伴う破断強度低下を補償する観点においては、前述のとおり、一定の減圧雰囲気での浸炭処理が好ましく、具体的には、該浸炭処理を、雰囲気圧力が３０ｈＰａ以下に調整された減圧浸炭雰囲気にて行なうことが望ましい。しかし、むやみに圧力を低下させると浸炭に要する時間が長くなり、製造コストの上昇を招くとともに、浸炭むらが発生しやすくなる。また、鋼中のＳｉは炭化物析出を抑制する作用と、酸化物膜の形成により浸炭を阻害する傾向とが強い元素である。この観点から、上記減圧浸炭雰囲気は、１ｈＰａ以上の圧力とすることが望ましい（特に、Ｓｉは浸炭阻害元素なので、その含有率が上限値である１．５質量％に近いものは、浸炭性を低下させないよう、上記以上の圧力で浸炭処理を行なうことが不可欠である）。このような範囲に浸炭処理時の雰囲気圧力を調整することで、浸炭処理における表面酸化や浸炭濃度むらや、浸炭深さ不足が生じにくくなり、かつ、面疲労強度に優れた浸炭焼入れ部材を実現することができる。
【００３９】
次に、部材の表面硬さを確保するには、焼入れ組織が十分に形成されることが重要である。焼入組織の主体をなすマルテンサイトが十分に生成するには、焼入冷却時において、Ｍｓ点近傍の温度を、一定の冷却速度（臨界冷却速度）以上で通過させなければならない。従って、焼入後の部材表面硬さを十分に高めるには、ガス冷却媒体の圧力を一定以上に高め、焼入れ組織を十分に形成するための冷却能を確保する必要がある。他方、冷却速度が過度に大きくなりすぎると（つまり、ガス冷却媒体の冷却能が高くなりすぎると）、焼入歪が大きくなり、良好な寸法及び形状精度が得られなくなる。従って、冷却能を決定するガス圧力には、ある適正な範囲が存在する。しかしながら、同じ冷却雰囲気であっても、部材の材質や寸法が異なる場合には、部材内部の冷却速度分布も変化し、必ずしも同じ焼入状態が得られるとは限らない。
【００４０】
この場合、ガス圧力に代えて焼入急冷度Ｈを用いると便利である。すなわち、部材を構成する鋼の熱伝導率をλ（単位：ｋｃａｌ／ｍｈ℃）、該ガス冷却雰囲気における部材の表面熱伝達係数をα（単位：ｋｃａｌ／ｍｈ２℃）として、焼入急冷度Ｈは、Ｈ≡０．５×（α／λ）にて定義される。このうち熱伝導率λは、鋼の材質により固有に決定される物性値であり、熱伝達係数αは、部材の比熱、熱伝導率、重量、形状及び寸法、さらに、ガス冷却媒体の種類（比熱）、圧力、流速等により決定され、冷却雰囲気と部材の形状、材質及び寸法が決まれば、周知の伝熱解析の主法により一義的に定まるパラメータである。そして、前記組成の鋼を採用する本発明にておいては、ガス焼入れを行なう際のガス冷却雰囲気を、部材の材質及び形状／寸法に応じて、ガス種、圧力及び流速等を調整し、上記焼入急冷度Ｈが０．０１〜０．０８となるように設定することが望ましい。Ｈが０．０１未満の冷却雰囲気では、焼入組織が十分に形成されず、部材の硬さが不足する。逆にＨが０．０８を超えると、焼入歪が大きくなり、良好な寸法及び形状精度が得られなくなる。
【００４１】
なお、ガス焼入れに使用する冷却ガスとしては、部材酸化抑制の観点から不活性ガス（例えばアルゴンガスなど）あるいは窒素ガスを用いて行なうことが望ましい。特に窒素ガスは、比熱が比較的大きく冷却能に優れ（１．０３Ｊ・Ｋ−１・ｇ−１、アルゴンは０．５２Ｊ・Ｋ−１・ｇ−１）、また、量産操業時における入手容易性とコスト及び取り扱い容易性などの点から本発明に好適に使用できる。
【００４２】
また、部材を構成する鋼としては、
Ｎ≡１０６×Ｃ（質量％）＋１０．８×Ｓｉ（質量％）＋１９．９×Ｍｎ（質量％）＋１６．７×Ｎｉ（質量％）＋８．５５×Ｃｒ（質量％）＋４５．５×Ｍｏ（質量％）＋２８
により表される成分パラメータＮが９５以下となるように組成調整された浸炭用鋼を使用するのがよい。Ｎが９５を超えると、鋼の圧延状態の硬さや焼ならし状態の硬さが著しく上昇し、機械加工性及び冷間加工性が得られなくなるからである。したがって、製造性を重視する場合にはこの成分パラメータＮが９５以下となるように鋼の成分組成を制御する必要がある。
【００４３】
また、浸炭処理後のガス焼入により、浸炭層の表面にて測定したビッカース硬度は７００Ｈｖ以上となるのがよい。浸炭焼入れ後の表面硬度は部材強度（特に疲労強度）に影響し、表面硬度は７００Ｈｖ未満では部材強度を十分に確保できなくなる場合がある。従って、特に疲労強度を重視する場合は、表面硬度を７００Ｈｖ以上とすることが望ましい。なお、浸炭層の表面にて測定したビッカース硬度の上限値に制限はなく、例えば９００Ｈｖ程度までならば少なくとも、炭化物析出等を抑制しつつ問題なく浸炭を行なうことができる。なお、表層におけるセメンタイト等の炭化物の生成が過剰となって表面硬度が９００Ｈｖを超える場合、かえって強度不足、特に靱性の低下が生ずる場合がある。
【００４４】
他方、浸炭層内側の非浸炭部のビッカース硬度は２５０Ｈｖ以上以下となるのがよい。浸炭層内側の非浸炭部のビッカース硬度が２５０Ｈｖ未満になると、内部起点の疲労破壊が起きやすくなり、疲労強度が低下する。したがって、特に疲労強度を重視する場合には、該非浸炭部のビッカース硬度を２５０Ｈｖ以上とするのがよく、これによって強度と靱性を合わせ持った部品が得られる。なお、本発明においてビッカース硬さは、ＪＩＳ：Ｚ２２４４（１９９８）に規定された試験方法により、試験荷重２．９４Ｎにて測定されたものをいう。
【００４５】
焼入れ後の非浸炭部の硬さを上記のように確保するには、ベイナイトが十分な量にて形成されている必要があり、望ましくは非浸炭部の組織がベイナイトを主体とするものになっているのがよい。なお、本明細書にて「ベイナイトが主体になる」とは、断面組織におけるベイナイトの面積率が５０％以上であることをいう。ベイナイトはマルテンサイトと異なり、格子をなす鉄原子が部分的に拡散しながら変態が進行する。従って、マルテンサイトと比較して変態に伴う歪みの発生が小さく、しかも、さらに冷却速度が小さくなったときに生成するパーライトよりは硬さが大きいので、内側の非浸炭部の強度を適度に高めることができる。内層部をベイナイトを主体に構成するには、部材寸法によっても異なるが、前述の限界冷却速度の測定において、ベイナイトが主体となる組織が得られる冷却速度が０．５℃／秒〜１０℃／秒の範囲に存在するように（より望ましくは、３℃／秒の冷却速度としたとき、ベイナイトが主体となる組織が得られるように）、組成選定することが望ましい。
【００４６】
次に、焼きならし状態の鋼素材の硬度は、前記した加工性改善の観点から、ロックウェルＢスケール硬さが９５ＨＲＢ以下であることが望ましい。なお、本発明においてロックウェルＢスケールは、ＪＩＳ：Ｚ２２４５（１９９８）に規定された試験方法により測定されたものをいう。
【００４７】
浸炭層表面の残留オーステナイト量は、２５％以下となっていることが望ましい。残留オーステナイト量が２５％を超えると硬さが低下し、特に重荷重あるいは衝撃等の加わりやすい歯車部材（例えば動力用や自動車用の歯車（例えば変速機用歯車）など）等への適用を図る場合、歯車の変形や歯面の波打ち（リップリング）といった初期損傷が発生しやすくなる問題がある。なお、浸炭層表面の残留オーステナイト量は、より望ましくは２０％以下であるのがよい。
【００４８】
他方、残留オーステナイトの面積率の下限に特に制限はなく、例えばショットピーニング等により、残留オーステナイトを強制的にマルテンサイト化して量を減らすこともできる。この場合、最終的な残留オーステナイト量を、例えば面積率にて、最大で１％（これは事実上、測定限界以下の値である）程度にまで減少させることもできる。なお、鋼のマルテンサイト変態は体積膨張が大きいので、その後背応力の影響を受けて残留するオーステナイトは、完全にはゼロにできない場合がある。
【００４９】
なお、残留オーステナイト量は、浸炭焼入れ層表面においてディフラクトメータ法によりＸ線回折プロファイルを測定したとき、体心立方晶系（フェライト）あるいは体心正方晶系（マルテンサイト）の相の主回折ピークの積分面積をＩｆ（｛２００｝及び｛２１１｝；正方晶系ではピークスプリットするので、グループに属する全てのピーク面積を合計する）と、面心立方晶系、すなわちオーステナイト相の回折ピークの積分面積をＩａ（｛２００｝、｛２２０｝、｛３１１｝｝として、
｛Ｉａ／（Ｉａ＋Ｉｆ）｝×１００（％）
にて表すものとする。
【００５０】
さらに、浸炭層表面のトルースタイト組織の面積率は、１０％以下であることが望ましい。トルースタイトは、浸炭焼入れ後の浸炭層に生成する不完全焼入組織であり硬さも小さい。従って、この浸炭層表面の面積率が１０％を超えると、面疲労強度が顕著に劣化する。そこで、特に面疲労強度を重視する場合には上記トルースタイトの面積率を１０％以下に規制することが望ましい。
【００５１】
また、部材表面からの粒界酸化が生じている深さは、３μｍ以内であることが望ましい。浸炭中に雰囲気から部材（鋼）に侵入する酸素は、粒界拡散が支配的となることが多く、粒界に酸化物相を形成する。粒界酸化物相が形成されると粒界強度が低下し、浸炭層の強度不足や脱粒等による耐摩耗性の低下を招く場合がある。また、粒界酸化物相の生成時に、周囲の鋼添加元素（副成分）の一部が粒界酸化物相に取り込まれる結果、粒界酸化物相周囲に添加元素の枯渇領域が生じ、浸炭焼入層の焼入性不足ひいては硬度や強度の不足を引き起こすおそれがある。従って、鋼材の組成調整、浸炭時の雰囲気（特に酸素分圧）、浸炭温度及び時間等を調整することにより、上記粒界酸化深さを３μｍ以下に抑制するようにする。
【００５２】
なお、粒界酸化相は部材断面を研磨することにより、非酸化領域と目視により簡単に判別できるので、その断面光学顕微鏡写真から上記の粒界酸化深さを測定することができる。
【００５３】
次に、部材の表面圧縮残留応力は３００〜８００ＭＰａとなっていることが望ましい。圧縮残留応力を部材表面に残留させることにより、表層部に亀裂が形成されたときに、その亀裂の拡大・伝播が抑制される。従って、部材の強度、特に動的強度（面疲労強度、曲げ疲労強度、ねじり疲労強度等）を大幅に向上させることができる。前記した通り、浸炭層に焼入れ処理してマルテンサイトを生成させると、変態に伴う体積膨張により圧縮応力場を生じ、上記のような表面圧縮残留応力を付与するのに好都合である。しかし、マルテンサイトの生成量が少ない場合、すなわち残留オーステナイトが多い場合は、十分な圧縮残留応力場を形成できない。従って、残留オーステナイトを減少させること（具体的には２５％以下とすること）は、このような圧縮残留応力効果を高める観点においても有利に作用するといえる。なお、マルテンサイト変態時の体積膨張の吸収は、マルテンサイト量が少ない場合は周囲のオーステナイトを塑性変形させて進行するため応力緩和し、表面圧縮残留応力の増大にはそれほど寄与しない。しかし、マルテンサイト量が増え残留オーステナイトが上記のように減少すると、塑性変形により導入された転位の密度が増加し、すべり変形が拘束されるため、表面圧縮残留応力は急速に増加する。また、焼入れ後にショットピーニング等の表層加工を施して圧縮残留応力を増加させる方法もある。後者の場合、ショットピーニング処理により残留オーステナイトをマルテンサイト化させると、圧縮残留応力を向上させる上でより有利となる。圧縮残留応力が３００ＭＰａ未満になると、亀裂伝播抑制による強度向上効果が十分得られなくなり、また、８００ＭＰａを超える圧縮残留応力を付与することは、マルテンサイト量を過度に多くするか、後加工による歪付与を過剰に大きくするかのいずれかを選択しなければならないが、前者は焼入れ時の冷却速度を、限度を超えて大きくしなければならず、後者は加工歪を過度に大きくしなければならなくなる。いずれも、部材の寸法精度確保という本発明の課題に照らし合わせれば、本末転倒の結果を招く。
【００５４】
なお、残留応力の測定は、浸炭焼入れ層表面においてディフラクトメータ法によりＸ線回折プロファイルを測定したとき、体心立方晶系（フェライト）あるいは体心正方晶系（マルテンサイト）の相の（２１１）ピークの半値幅とピーク中心位置との関係に基づいて行なうことができる。この方法により応力分析を行なう装置は、Ｆａｓｔｒｅｓｓ（登録商標名）応力分析器として周知であり（例えば「新版Ｘ線回折要論」（カリティ著、アグネ（１９７９）、４３１〜４３３頁）、詳細な説明は省略する。
【００５５】
本発明の浸炭焼入れ部材は、上記本発明の製造方法により実現可能なものであり、部材非浸炭部を構成する鋼が：副成分の含有量が、Ｃ：０．１２〜０．２２質量％、Ｓｉ：０．４〜１．５質量％、Ｍｎ：０．２５〜０．４５質量％、Ｎｉ：０．５〜１．５質量％、Ｃｒ：１．３〜２．３質量％、Ｂ：０．００１〜０．００３質量％、Ｔｉ：０．０２〜０．０６質量％、Ｎｂ：０．０２〜０．１２質量％、Ａｌ：０．００５〜０．０５質量％、Ｍｏ：０．０５質量％以下に設定され、残部が主成分をなすＦｅと不可避不純物に設定され；
Ｎ≡１０６×Ｃ（質量％）＋１０．８×Ｓｉ（質量％）＋１９．９×Ｍｎ（質量％）＋１６．７×Ｎｉ（質量％）＋８．５５×Ｃｒ（質量％）＋４５．５×Ｍｏ（質量％）＋２８により表される成分パラメータＮが９５以下となり；
また、９００℃から室温まで一定速度にて冷却したとき、冷却速度が少なくとも０．１℃／秒以下の範囲においてはベイナイトが生成せず、かつ、冷却速度が少なくとも１２℃／秒以上の範囲においてはフェライトが生成しなくなるように；
前記副成分含有量が調整された鋼であり、前記部材表面に浸炭層が形成されるとともに、焼入れ後の該浸炭層の表面にて測定したビッカース硬度が７００Ｈｖ以上であり、また、焼入れ後の浸炭層内側の非浸炭部のビッカース硬度が２５０Ｈｖ以上であることを特徴とする。
【００５６】
上記本発明の浸炭焼入れ部材は、上記組成の鋼の採用により、浸炭焼入前の加工が容易であり、また、部材寸法に関係なくガス冷却により十分な焼入組織が形成可能であるから、低歪で高精度の部材が実現できる。また、浸炭焼入れ後の破断強度や面疲労強度に優れ、さらに部材全体の製造コストを大幅に削減できる。各数値の臨界的意味は、すでに説明済みであるから省略する。なお、浸炭層の組成は、非浸炭部の組成をベースとして、浸炭により炭素含有量を増加させたものに相当する。
【００５７】
【実施例】
以下、本発明の効果を確認するために行なった実験結果について説明する。
まず、表１に示す化学組成の鋼をアーク炉で溶製後、熱間圧延により直径１５０ｍｍ及び直径３２ｍｍの丸棒とし、９２５℃に１時間保持後空冷の焼ならしを行った。鋼種Ａ１、Ａ２、Ａ３は本発明の請求項に記載されている鋼材組成に該当する鋼種であり、鋼種Ｂ及び鋼種ＣはそれぞれＪＩＳの肌焼鋼ＳＣＭ４２０及びＳＮＣＭ８１５に相当する鋼種である（成分パラメータＮの計算値も合わせて示している）。
【００５８】
【表１】
すべての鋼種について、直径３２ｍｍの焼ならし材の、横断面の中心部についてロックウェル硬さを測定した。そして、各鋼材のベイナイト生成の下限臨界冷却速度βLC及びフェライト生成の上限臨界冷却速度βUCを、図４の試験片を別途作成することにより、すでに説明した方法にて測定した。また、直径２５ｍｍ、長さ５０ｍｍの丸棒試験片及び図１に示す形状の回転曲げ疲れ試験片を加工した。また、直径３２ｍｍ及び直径１５０ｍｍの焼ならし材から、図２に示すピッティング試験用ローラー及びその相手ローラーを加工した。他方、直径１５０ｍｍの焼ならし材は、軸直交断面により切断し、さらに鍛造及び切削により、図３に示す形状の歯車を加工した。
【００６０】
これらの部材のうち、直径２５ｍｍの丸棒試験片については、鋼種Ａ１、Ａ２、Ａ３、Ｂ、Ｃとも、表２に示す条件で低圧浸炭及びガス焼入れを行った（それぞれ発明例Ａ、比較例Ｂ、比較例Ｃとする）。そして、歯車、回転曲げ疲れ試験片、ピッティング試験ローラー及び相手ローラーについては、鋼種Ａ１、Ａ２、Ａ３及び鋼種Ｃは表２に示した条件で低圧浸炭及びガス焼入れを行い（発明例Ａ１、Ａ２、Ａ３、比較例Ｃ）、鋼種Ｂは表３に示す条件でガス浸炭及び油焼入れを行った。
【００６１】
【表２】
【表３】
また、試験片の軸断面を鏡面研磨し、表層部のＸＭＡ（X-ray Micro Analysis）分析を行なうことにより、浸炭後の表層部の炭素濃度を、表面から５０μｍの位置にて調べた（結果を表１に示している）。他方、直径２５ｍｍの丸棒試験片について、ビッカース硬度計により横断面の硬さ分布を調べた。なお、浸炭焼入材の表層硬さは、表面から０．０２ｍｍの位置において測定した。さらに、これと同等の位置においてトルースタイトの面積率を、走査型電子顕微鏡写真を画像解析することにより測定した。また、すでに説明した方法により、表層部の残留オーステナイト量、圧縮残留応力を測定した。これらの結果を表４に示す。
【００６４】
【表４】
表４において、発明例Ａ１〜Ａ３は９５ＨＲＢ以下の焼ならし硬さを示し、浸炭焼入れ材の中心部の硬さは３５０Ｈｖ以上である。表層（浸炭層）の組織はいずれもマルテンサイトであり、中心部（非浸炭部）の組織はいずれもベイナイトあるいはベイナイト＋マルテンサイト（ただし、ベイナイトが５０％以上）であり、顕著な不完全焼入組織は存在していなかった。これに対し、比較例Ｂは焼ならし硬さは低いものの、浸炭焼入れ材の表層硬さ及び中心部硬さが発明例に対して低い。また、比較例Ｃは浸炭焼入れ材の表層硬さ及び中心部硬さは発明例Ａとほぼ同等であるが、焼ならし硬さが極めて高い。
【００６６】
回転曲げ疲れ試験は、小野式回転曲げ疲れ試験機を用い、繰り返し数１千万回を基準とする疲労強度を求めた。また、ピッティング試験は潤滑油中で試験ローラーと相手ローラーを荷重とすべり速度を制御しつつ回転接触させ、試験ローラーの表面に発生するピッティング損傷を振動により検出し、繰り返し数１千万回を基準とする面圧強度を求めた。なお、相対すべり速度は８００ｍｍ／ｓとした。表４にこれらの強度試験の結果を示す（ただし、各強度はいずれも比較例Ｂを基準とする相対値にて表している）。この結果からわかるように、発明例では回転曲げ疲労強度及び耐ピッティング強度において比較例を大きく上回る特性が得られている。
【００６７】
次に、歯車については、浸炭焼入れ後の精度を、以下の方法により評価した。まず、歯車精度については、専用の精密ギヤ精度測定機に歯車をセットし、左右歯面それぞれにおいて、歯車の圧力角方向の誤差量とネジレ角方向の誤差量を測定した。また、歯溝の高さを全周測定し最大値から最小値を差し引いた値を歯溝の振れとして算出した。他方、歯車の寸法精度については、以下のようにして測定を行った。すなわち、歯車の互いにむかいあった２つの歯溝にボールを入れ、その外周寸法を専用のＯ．Ｂ．Ｄ（Over Ball Diaphragm）測定器にて測定した。Ｏ．Ｂ．Ｄの測定位置は図３に示すように、円周方向は直角２方向（Ｘ，Ｙ）であり、歯幅方向は上・中・下の３個所（Ａ，Ｂ，Ｃ）とした。表中「Ｏ．Ｂ．Ｄ楕円」とあるのは直角２方向でのＯ．Ｂ．Ｄの差の絶対値であり、「Ｏ．Ｂ．Ｄテーパー」は歯幅方向での上部Ｏ．Ｂ．Ｄと下部Ｏ．Ｂ．Ｄの差である。以上の結果を表５に示す。
【００６８】
【表５】
すなわち、本発明による浸炭焼入れ部材として構成した歯車は、熱処理歪が極めて小さいため、歯車精度及び寸法精度の全てにわたって非常に良好な特性が得られていることがわかる。
【図面の簡単な説明】
【図１】回転曲げ疲れ試験片を示す図。
【図２】ピッティング試験用ローラー及び相手ローラーを示す図。
【図３】熱処理歪評価用歯車を示す図。
【図４】連続冷却変態線図測定用の試料の形状を示す図。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for manufacturing a carburized and quenched member excellent in fatigue strength and shape accuracy, and a carburized and quenched member.
[0002]
[Prior art]
Carburizing and quenching is a heat treatment in which quenching is performed after increasing the carbon concentration of the surface layer of steel. Since the carburized and quenched steel parts have high surface hardness and high internal toughness, they are applied to power transmission mechanism parts that require fatigue strength and toughness. In addition, since the carbon content of the material before carburizing and quenching is relatively small, machine removal processing such as cutting and polishing, and plastic processing such as rolling or forging are relatively easy, and it is excellent in part manufacturability, so mass production It is widely used in parts such as automobiles. Accordingly, carburizing and quenching steel is required on the premise that the manufacturing cost of parts is low as well as strength.
[0003]
Since the surface layer of the carburized and quenched member has the highest hardness as steel, if there are minute defects or fragile parts in the vicinity of the surface, these act as starting points for fatigue failure or impact failure, so the strength of the member is high. Leading to a decline. Moreover, since a general carburizing process is performed at a high temperature in an oxidizing metamorphic gas atmosphere, the surface of the member to be processed is easily oxidized. Such surface oxidation is known to cause a decrease in hardenability of the surface layer portion due to depletion of alloy elements and to deteriorate member strength. As a method for improving this, vacuum carburizing treatment or applying high-strength steel to which non-oxidizing elements are added while reducing elements that promote oxidation has been performed. In addition, when wear resistance or pitting resistance is required, it is described in Japanese Patent Publication No. 7-116551 that the addition of an element that improves high-temperature strength and softening resistance is effective.
[0004]
[Problems to be solved by the invention]
The strength and cost required for the above carburized and quenched material are not always sufficient. In particular, it is the most important issue in carburized and quenched parts that both fatigue strength and impact strength in strength and wear resistance and pitting resistance are compatible, and that both strength and manufacturing cost such as machinability and shape accuracy are compatible. However, the conventional technology does not sufficiently meet today's market demand for weight reduction and cost reduction. The reason will be described below.
[0005]
Machining accounts for the majority of the manufacturing cost of carburized and quenched members. The material hardness has the greatest influence on the machinability before carburizing and quenching. For example, the high-strength steel described in the above-mentioned Japanese Patent Publication No. 7-116551 or the like is generally a high alloy steel and is therefore inferior in machinability.
[0006]
For example, regarding fatigue strength and impact strength (hereinafter collectively referred to as “breaking strength”) and wear resistance and pitting resistance (hereinafter collectively referred to as “surface fatigue strength”) If each is independent, it can be improved by adjusting the chemical composition of the steel material. For example, it is known that the reduction of Si, Mn and Cr and the addition of Ni and Mo are effective for the breaking strength, and the addition of Si, Cr and Mo is effective for the surface fatigue strength. . However, such high alloying increases the cost of machining due to the increase in material hardness, in addition to the increase in steel material cost, or causes deterioration of carburizing property and increase in retained austenite structure during carburizing and quenching, There is a limit.
[0007]
On the other hand, among chemical components that affect the strength of the carburized and quenched member, Si is the element that is most harmful to the fracture strength, and the most useful element for the surface fatigue strength. Therefore, in order to achieve both fracture strength and surface fatigue strength as required by gears and the like only with chemical components of steel, a large amount of relatively expensive strengthening elements other than Si, such as Ni and Mo, are used. It is necessary to add, and as described above, the adverse effect of high alloying becomes large, so that both strength and manufacturability cannot be achieved.
[0008]
Next, when the required shape and dimensional tolerances of parts are severe, if the heat treatment strain generated during carburizing and quenching is large, the required dimensional accuracy cannot be ensured in the state after carburizing and quenching, and grinding finish is required. There is. Especially in gears and the like, since the shape accuracy causes noise and damage, tooth surface grinding may be performed by a dedicated processing machine. Such grinding is generally more costly than cutting and sometimes accounts for 50% or more of the overall manufacturing cost of the carburized and quenched member. Therefore, at the time of carburizing and quenching, the generated strain is required to be as small as possible. Conventionally, improvement has been achieved mainly by quenching technology, but has not been sufficiently improved.
[0009]
In most cases, oil or salt is used as a cooling medium in carburizing and quenching, but quenching by gas cooling (hereinafter referred to as “gas quenching”) may also be performed. Gas quenching is advantageous from the viewpoint of reducing the processing cost by ensuring the accuracy of dimensions and shape because the heat treatment distortion is small. However, since the gas has a low cooling capacity, the quenching structure is not sufficiently formed and may be incompletely quenched. In such an incompletely hardened state, sufficient strength and mechanical properties cannot be obtained, and as a result, gas cooling can only be applied to small parts.
[0010]
For example, increasing the gas pressure in the cooling chamber is effective for improving the cooling capacity of the gas quenching, but increases the size of the equipment and increases the processing cost. For such a lack of cooling ability, it is effective to increase the hardenability of the steel material, but typical elements such as Si, Mn, Cr, Ni, and Mo that enhance the hardenability of the steel are carburized. It is also an element that increases the hardness of the material before quenching, is harmful to machinability and cold workability, and deteriorates the manufacturability of parts. Moreover, in JIS (Japanese Industrial Standard), SNCM616, SNCM815, etc. are prescribed | regulated as case hardening steel with high hardenability, However, These are compared with SCr420 and SCM420 which are the most common case hardening steel, Due to the high alloy content, the hardness in the normalized and annealed state is very high, and cutting and cold working are difficult. That is, it was essentially difficult to achieve both the hardenability suitable for gas quenching and the manufacturability of parts suitable for mass production.
[0011]
The problem of the present invention is that it is easy to process before carburizing and quenching, and can perform carburizing and quenching with low distortion and high accuracy by gas cooling regardless of the member dimensions. An object of the present invention is to provide a method of manufacturing a carburized and quenched member that is excellent in strength and can significantly reduce the manufacturing cost of the entire member, and a carburized and quenched member obtained thereby.
[0012]
[Means for solving the problems and actions / effects]
In order to solve the above-mentioned problems, the method for producing a carburized and quenched member according to the present invention hasContent of, C: 0.12-0.22 mass%, Si: 0.4-1.5 mass%, Mn: 0.25-0.45 mass%, Ni: 0.5-1.5 mass%, Cr : 1.3 to 2.3 mass%, B: 0.001 to 0.003 mass%, Ti: 0.02 to 0.06 mass%, Nb: 0.02 to 0.12 mass%, Al: 0 0.005 to 0.05 mass%,Mo: set to 0.05% by mass or less, the balance is set to Fe and inevitable impurities as the main component,In addition, when cooling at a constant rate from 900 ° C. to room temperature, bainite is not generated in the range where the cooling rate is at least 0.1 ° C./second or less, and in the range where the cooling rate is at least 12 ° C./second or more. Constitutes a member with steel in which the content of each subcomponent in the non-carburized state is adjusted so that ferrite is not generated,In a reduced pressure atmosphere of 1-30 HPaBy carburizing the member, a carburized layer having a surface carbon concentration of 0.6 to 1.5% by mass is formed on the member, and after the carburizing treatment, the member isWhile quenching by gas cooling using an inert gas or nitrogen gas so that the quenching quenching degree H becomes 0.01 to 0.08,
In the gas cooling atmosphere when performing the gas quenching, the thermal conductivity of the steel constituting the member is λ (unit: kcal / mh ° C.), and the surface heat transfer coefficient of the member in the gas cooling atmosphere is α (unit: kcal / mh2 ° C.), an atmosphere in which the quenching quenching degree H defined by H≡0.5 × (α / λ) is 0.01 to 0.08 is used.It is characterized by that.
[0013]
According to the method of the present invention, processing before carburizing and quenching is easy, and high-precision carburizing and quenching can be performed with low distortion by gas cooling regardless of member dimensions. The surface fatigue strength is excellent, and the manufacturing cost of the entire member can be greatly reduced. The reason will be described in detail below.
[0014]
The inventors of the present invention studied the chemical composition and heat treatment of steel materials, including their interactions, with the aim of simultaneously improving all of the breaking strength, surface fatigue strength, material hardness and heat treatment strain of carburized and quenched materials. As a result, the following was found. First, the use of Si is indispensable for improving the surface fatigue strength, but excessive addition has an extremely harmful effect on the breaking strength. On the other hand, in the carburized and hardened member, when the surface of the member to be treated is oxidized, the hardenability of the surface layer portion is reduced due to the depletion of the alloy elements, and similarly the strength of the fracture is deteriorated. Therefore, by carburizing the member in a reduced pressure carburizing atmosphere or a carburizing atmosphere mainly composed of an inert gas, the effect of surface oxidation is greatly reduced, and a margin for a decrease in fracture strength caused by adding Si is secured. Can do.
[0015]
Next, in order to apply gas quenching effective for reducing heat treatment strain to members having a wide range of shapes and sizes, it is effective to improve the hardenability of the steel material. In the present invention, in consideration of adopting a reduced pressure carburizing atmosphere or a carburizing atmosphere mainly composed of an inert gas, when cooling at a constant rate from 900 ° C. to room temperature, the cooling rate is at least 0.1 ° C./second. In the following range, steel that does not generate bainite and that does not generate ferrite in a region where the cooling rate is at least 12 ° C./second or more is employed.
[0016]
As the cooling rate increases, the steel structure is generally known to change in the following order:
F + P → (F + P + B → F + B → B → B + M) → M
Here, F represents ferrite, P represents pearlite, B represents bainite, and M represents martensite. Depending on the alloy composition, there may be no conditions for single bainite. Pearlite is a eutectoid structure of ferrite and cementite, and the formation of pearlite also means that ferrite is generated. F + P, F + B, and the like mean that there is ferrite formed independently in a form other than pearlite. The range of the cooling rate can be specified by measuring a continuous cooling transformation diagram (CCT diagram) of steel at various cooling rates.
[0017]
As described above, martensite plays a key role in quench hardening of the carburized layer. As is well known, the martensitic transformation is caused by a coordinated movement of the crystal lattice in a pseudo-shear deformation without large diffusion of atoms. Since the martensitic transformation of steel has a large volume expansion, it proceeds in a form that greatly distorts the surrounding retained austenite. This also contributes to quench hardening of steel. However, the fact that the strain associated with the transformation is large means that the hardness is improved, but it is also difficult to cause a dimensional deviation in the member after quenching.
[0018]
In the case of the present invention, a carburized layer is formed on the surface of the member, and martensite is preferentially generated in the carburized layer, thereby increasing the hardness of the surface layer portion and improving the wear resistance. In the present invention, gas cooling, which has a cooling rate lower than that of water or oil, is employed in order to suppress an increase in distortion due to cooling. Therefore, the composition of the steel is set so that ferrite is not generated in the region where the cooling rate is at least 12 ° C./second (hereinafter referred to as the upper limit cooling rate) so that the carburized layer can be sufficiently quenched even with gas cooling. To do. If ferrite is generated even if the cooling rate is increased to 12 ° C./second or more, martensite is not sufficiently formed in the carburized layer by gas cooling, leading to insufficient hardness.
[0019]
However, the hardenability becomes excessively good, which is disadvantageous in the present invention. That is, if the martensite is excessively generated even in the inner layer portion that is not affected by carburization, the amount of martensite generated as a whole member increases, leading to a decrease in dimensional accuracy. Therefore, it is important to select the composition so that martensite is sufficiently generated in the carburized layer during gas quenching, but excessive martensite is not formed in the inner layer portion. Specifically, bainite is prevented from being generated in a range where the cooling rate is at least 0.1 ° C./second or less (hereinafter referred to as the lower limit cooling rate). If bainite is generated even at a cooling rate of 0.1 ° C./second or less, the inner layer portion that is not affected by the carburized layer is deeply baked and distortion increases, and the object of the present invention cannot be achieved.
[0020]
On the other hand, another object of the present invention is to enable easy processing before carburizing and quenching, and for that purpose, the annealing state, that is, the range in which the cooling rate from austenite corresponds to cooling or air cooling. In the following (hereinafter referred to as annealing and cooling rate range), a material hardness sufficiently low to improve workability must be obtained. In view of this, when the cooling rate is set to be less than 0.1 ° C./second, the composition of the non-carburized steel is selected so that bainite is not generated. Therefore, if the bainite is not generated at a cooling rate of less than 0.1 ° C./second as described above, the generation of bainite is sufficiently suppressed in the range of the actual annealing cooling rate, and the workability of the ferrite + pearlite is large. A rich organization can be obtained.
[0021]
In addition, as a method for experimentally determining the critical cooling rate for generating ferrite or bainite in the inner layer portion, after normalizing the steel, it is processed into a test piece having the shape and dimensions shown in FIG. A bottomed hole is formed in this test piece from one end to the center in the longitudinal direction, and a thermocouple is fixed to the bottom of the hole by welding. In this state, the test piece is placed in a predetermined constant flow rate of cooling gas, and further heated to 900 ° C., which is the austenitizing temperature, by high frequency induction heating. If the temperature is kept substantially constant at 900 ° C., the temperature at the center position of the test piece is monitored by a thermocouple while the power of high-frequency induction heating is controlled to be reduced, and cooling is performed at a constant cooling rate. During cooling, the test piece is fixed at one end, and the position of the support rod attached to the other end is monitored using a differential transformer. Read the temperature of ferrite or bainite formation from the inflection point. Further, after the cooling is completed, the test piece is cut along the axis orthogonal plane at a position corresponding to the thermocouple, and the presence or absence of ferrite or bainite is visually confirmed by observation with an optical microscope. This measurement is performed while changing the cooling rate in various ways, and the critical cooling rate is determined by checking the presence or absence of ferrite and bainite. In this measurement method, the formation of ferrite and bainite is confirmed by both the inflection point of the test piece dimensional change and the optical microscope structure. In the former case, the inflection point is determined by a well-known differential analysis. If there is an error and no inflection point corresponding to the transformation is found within the range of the error, it is determined that no ferrite or bainite is generated. On the other hand, when observing with an optical microscope structure, nine fields of view are randomly selected from the 0.1 mm square on the cross section, and no ferrite or bainite is generated unless ferrite or bainite is visually observed in each field of view. Is determined. In either case, even if ferrite or bainite of about 1 to 2% in terms of area ratio is generated, it is determined that there is substantially no generation of ferrite or bainite from the problem of measurement limit.
[0022]
As described above, in order to obtain hardenability suitable for gas quenching from the viewpoint of both heat treatment strain and strength, it is effective to increase the amount of alloying elements, but it is necessary to minimize the adverse effects of high alloying. is there. On the other hand, in order to obtain a low material hardness, it is necessary to suppress the formation of bainite and to have a structure with a large amount of ferrite and pearlite in a range where the cooling rate from austenite corresponds to cooling or air cooling. In general, it has been difficult to achieve both high hardness after quenching and low hardness of the material. However, in the present invention, it is possible to achieve both in a preferable range by combining specific alloy elements. This will be described in detail below.
[0023]
First, boron (B) is the most effective element as a chemical component that produces the above effects. In the case where B is present in a solid solution state in austenite, the cooling rate is 2 ° C./second or more, and the formation of ferrite and bainite is suppressed and the hardenability is improved. However, when the cooling rate is 1 ° C./second or less, The same structure and hardness as when B is not added can be obtained.
[0024]
On the other hand, adjustment of hardenability and normalization hardness is made by adding Si, Mn, Cr, Ni, Mo, and the like. Here, in general case-hardened steel and the like, the behavior of these contents has been analyzed sufficiently and a certain degree of prediction has been possible. However, in the present invention, addition of B is essential in order to obtain a low material hardness, and solid solution B always exists in the steel material. Regarding the steel having such a solid solution B, the influence of the alloy elements, that is, the interaction with B, has not been sufficiently clarified heretofore, and naturally there has been no proposal for their proper composition range. .
[0025]
As a result of detailed studies on this point, the present inventors have found that Ni and Si have a particularly large contribution to the improvement of hardenability when coexisting with solute B, and conversely, the influence on the material hardness is small. I found it. Conversely, it has also been found that Mn and Mo have a particularly large effect of increasing the material hardness when coexisting with solute B. The steel composition adopted in the present invention is to achieve both high hardenability and low material hardness by adjusting the alloy element content based on the addition or increase of B, Ni, Si, and reduction of Mn, Mo. It is possible and unique.
[0026]
Hereinafter, the reasons for limiting the composition of steel employed in the present invention will be described.
(1) C: 0.12-0.22 mass%
C is an element that improves the strength of the non-carburized portion of the carburized and quenched material. However, if the content is less than 0.12% by mass, the effect is small, and if it exceeds 0.22% by mass, the material hardness increases. Therefore, the content of C is set to 0.12 to 0.22% by mass.
[0027]
(2) Si: 0.40 to 1.50 mass%
Si is an element effective for improving the surface fatigue strength, and also improves the hardenability especially when coexisting with the solute B and has a relatively small influence on the material hardness. It is an element to be added. If the content is less than 0.40% by mass, the effect of improving the surface fatigue strength and hardenability is small, and if it exceeds 1.50% by mass, it becomes difficult to form uniform austenite before quenching due to an increase in the A3 transformation point. Therefore, the Si content is set to 0.40 to 1.50 mass%.
[0028]
(3) Mn: 0.25 to 0.45% by mass
Mn is an element that improves hardenability, but the effect is small if it is less than 0.25% by mass, and the effect of increasing the material hardness when coexisting with solute B is more than 0.45% by mass, It significantly degrades part manufacturability Therefore, the Mn content is set to 0.25 to 0.45 mass%.
[0029]
(4) Ni: 0.50 to 1.50 mass%
Ni is an element that improves the strength and hardenability of carburized steel, and its effect is particularly great when it coexists with solute B. However, when the content is less than 0.50% by mass, the effect is not remarkable. When the content exceeds 1.50% by mass, a bainite structure is likely to be generated, and the machinability is remarkably deteriorated. Therefore, the Ni content is 0.50 to 1.50 mass%.
[0030]
(5) Cr: 1.30-2.30% by mass
Cr is an element that improves carburization, and is an element that has an effect of preventing deterioration of carburization due to Si and Ni. However, if the content is less than 1.30% by mass, the effect is insufficient, and if the content exceeds 2.30% by mass, a bainite structure is easily formed in the material and the machinability is remarkably deteriorated. Therefore, the Cr content is 1.30 to 2.30 mass%.
[0031]
(6) B: 0.0010 to 0.0030 mass%
B is an element that remarkably improves hardenability, but if it is less than 0.0010% by mass, a stable effect cannot be obtained, and even if added over 0.0030% by mass, the effect is saturated, so it is economical. Not. Therefore, the content of B is set to 0.0010 to 0.0030% by mass.
[0032]
(7) Ti: 0.02 to 0.06% by mass
Ti is an element that prevents nitrides from becoming nitrides by forming nitrides and stabilizes the effect of improving hardenability by B. However, if the content is less than 0.02% by mass, the effect is small, and even if added over 0.06% by mass, the effect is saturated, and thus it is not economical. Therefore, the Ti content is 0.02 to 0.06 mass%.
[0033]
(8) Nb: 0.02 to 0.12% by mass
Nb is an element effective in suppressing the growth of crystal grains and improving the strength. If the content is less than 0.02% by mass, the effect is small, and if it exceeds 0.12% by mass, coarse carbonitrides are formed at the time of solidification and the effect of suppressing crystal grain growth is reduced. Therefore, the Nb content is set to 0.02 to 0.12% by mass.
[0034]
(9) Al: 0.005 to 0.050 mass%
Al promotes deoxidation in the melting process and is effective in reducing the amount of oxide inclusions, but if it is less than 0.005% by mass, there is no deoxidation effect, and it is added in excess of 0.050% by mass. Is not economical because the effect is saturated. Therefore, the content of Al is set to 0.005 to 0.050 mass%.
[0035]
Mo, which is often added to general carburizing steel, is an element that has a particularly large effect of improving hardenability when coexisting with B. However, since the material hardness in the state of hot working or normalization is significantly increased, in the present invention for the purpose of achieving both manufacturability and strength, it is not an element to be actively added, preferably, It is preferable to limit the Mo content in the steel to be employed to 0.05% by mass or less.
[0036]
Furthermore, in this invention, the surface carbon concentration of a carburized layer shall be 0.6-1.5 mass%. The surface carbon concentration of the carburized layer affects the surface hardness of the carburized quenching material, but the surface hardness is insufficient if it is less than 0.6% by mass, and the amount of carbide precipitated increases if it exceeds 1.5% by mass, The hardenability of the base is significantly reduced, leading to insufficient surface hardness.
[0037]
In addition, nitrogen gas is inexpensive in the carburizing atmosphere and can be suitably used in the present invention, but an inert gas such as argon can be used instead of the nitrogen gas.
[0038]
Hereinafter, requirements that can be further added in the method for manufacturing a carburized and quenched member of the present invention will be described. In the present inventors, steel containing up to 1.5% by mass of Si is adopted. However, from the viewpoint of compensating for the decrease in fracture strength accompanying silicon addition, as described above, carburizing treatment in a constant reduced-pressure atmosphere is performed. More specifically, it is desirable to perform the carburizing treatment in a reduced pressure carburizing atmosphere in which the atmospheric pressure is adjusted to 30 hPa or less. However, if the pressure is reduced unnecessarily, the time required for carburizing becomes longer, resulting in an increase in manufacturing costs and carburizing unevenness. Further, Si in steel is an element having a strong action of suppressing carbide precipitation and a tendency to inhibit carburization by forming an oxide film. From this point of view, it is desirable that the reduced pressure carburizing atmosphere has a pressure of 1 hPa or more (particularly, since Si is a carburization-inhibiting element, the one whose content is close to the upper limit of 1.5% by mass is carburizing. Carburizing treatment is indispensable at a pressure higher than the above so as not to decrease.) By adjusting the atmospheric pressure during carburizing treatment to such a range, carburizing and quenching members with excellent surface fatigue strength and surface oxidation, uneven carburizing concentration unevenness, and insufficient carburizing depth are realized. can do.
[0039]
Next, in order to ensure the surface hardness of the member, it is important that the quenched structure is sufficiently formed. In order to sufficiently generate martensite forming the main body of the quenched structure, the temperature in the vicinity of the Ms point must be passed at a constant cooling rate (critical cooling rate) or more during quenching cooling. Therefore, in order to sufficiently increase the hardness of the surface of the member after quenching, it is necessary to increase the pressure of the gas cooling medium to a certain level or more and ensure cooling ability for sufficiently forming a quenched structure. On the other hand, if the cooling rate is excessively high (that is, if the cooling capacity of the gas cooling medium is excessively high), the quenching distortion increases, and good dimensional and shape accuracy cannot be obtained. Therefore, there is a certain appropriate range for the gas pressure that determines the cooling capacity. However, even in the same cooling atmosphere, if the material and dimensions of the members are different, the cooling rate distribution inside the members also changes, and the same quenching state is not always obtained.
[0040]
In this case, it is convenient to use the quenching quenching degree H instead of the gas pressure. That is, the thermal conductivity of steel constituting the member is λ (unit: kcal / mh ° C.), and the surface heat transfer coefficient of the member in the gas cooling atmosphere is α (unit: kcal / mh).2The quenching quenching degree H is defined as H≡0.5 × (α / λ). Among these, the thermal conductivity λ is a physical property value uniquely determined by the material of the steel, and the heat transfer coefficient α is the specific heat of the member, the thermal conductivity, the weight, the shape and dimensions, and the type of the gas cooling medium ( Specific heat), pressure, flow rate, and the like. If the cooling atmosphere and the shape, material, and dimensions of the member are determined, the parameters are uniquely determined by a well-known main method of heat transfer analysis. And in the present invention employing the steel of the above composition, the gas cooling atmosphere at the time of gas quenching is adjusted according to the material and shape / dimension of the member, the gas type, pressure, flow rate, etc. It is desirable to set the quenching quenching degree H to be 0.01 to 0.08. In a cooling atmosphere where H is less than 0.01, a hardened structure is not sufficiently formed, and the hardness of the member is insufficient. On the other hand, if H exceeds 0.08, the quenching distortion increases, and good dimensional and shape accuracy cannot be obtained.
[0041]
In addition, as a cooling gas used for gas quenching, it is desirable to perform using inert gas (for example, argon gas etc.) or nitrogen gas from a viewpoint of member oxidation suppression. Nitrogen gas in particular has a relatively large specific heat and excellent cooling ability (1.03 J · K-1・ G-1Argon is 0.52 J · K-1・ G-1In addition, it can be suitably used in the present invention from the viewpoints of availability, cost, and ease of handling during mass production operations.
[0042]
Moreover, as steel constituting the member,
N≡106 × C (mass%) + 10.8 × Si (mass%) + 19.9 × Mn (mass%) + 16.7 × Ni (mass%) + 8.55 × Cr (mass%) + 45.5 × Mo (Mass%) +28
It is preferable to use carburizing steel whose composition is adjusted so that the component parameter N represented by This is because when N exceeds 95, the hardness in the rolled state and the hardness in the normalized state of the steel are remarkably increased, and the machinability and the cold workability cannot be obtained. Therefore, when emphasizing manufacturability, it is necessary to control the component composition of steel so that the component parameter N is 95 or less.
[0043]
Further, the Vickers hardness measured on the surface of the carburized layer by gas quenching after the carburizing treatment is preferably 700 Hv or more. The surface hardness after carburizing and quenching affects the member strength (particularly fatigue strength), and if the surface hardness is less than 700 Hv, the member strength may not be sufficiently secured. Therefore, it is desirable that the surface hardness be 700 Hv or more, particularly when fatigue strength is important. In addition, there is no restriction | limiting in the upper limit of the Vickers hardness measured on the surface of the carburized layer, and if it is up to about 900 Hv, for example, carburization can be performed without any problem while suppressing carbide precipitation. In addition, when generation | occurrence | production of carbide | carbonized_materials, such as cementite, in a surface layer becomes excessive, and surface hardness exceeds 900 Hv, intensity | strength lack, especially a toughness fall may arise.
[0044]
On the other hand, the Vickers hardness of the non-carburized portion inside the carburized layer is preferably 250 Hv or more. When the Vickers hardness of the non-carburized portion inside the carburized layer is less than 250 Hv, fatigue failure at the internal origin tends to occur, and the fatigue strength decreases. Therefore, when emphasizing fatigue strength in particular, the Vickers hardness of the non-carburized portion is preferably set to 250 Hv or more, thereby obtaining a component having both strength and toughness. In the present invention, the Vickers hardness means a value measured at a test load of 2.94 N by a test method defined in JIS: Z2244 (1998).
[0045]
In order to ensure the hardness of the non-carburized part after quenching as described above, the bainite needs to be formed in a sufficient amount, and the structure of the non-carburized part is preferably composed mainly of bainite. It is good to have. In the present specification, “being composed mainly of bainite” means that the area ratio of bainite in the cross-sectional structure is 50% or more. Unlike martensite, bainite undergoes transformation while the latticed iron atoms partially diffuse. Therefore, compared with martensite, the occurrence of distortion due to transformation is small, and since the hardness is higher than the pearlite produced when the cooling rate is further reduced, the strength of the inner non-carburized portion is appropriately increased. be able to. Although the inner layer portion is mainly composed of bainite, the cooling rate at which the structure mainly composed of bainite is obtained is 0.5 ° C./second to 10 ° C./second in the above-described measurement of the critical cooling rate. It is desirable to select the composition so that it exists in the range of seconds (more desirably, when a cooling rate of 3 ° C./second is obtained, a structure mainly composed of bainite is obtained).
[0046]
Next, the hardness of the normalized steel material is such that Rockwell B scale hardness is 95H from the viewpoint of improving the workability described above.RBThe following is desirable. In the present invention, the Rockwell B scale is measured by a test method defined in JIS: Z2245 (1998).
[0047]
The amount of retained austenite on the surface of the carburized layer is desirably 25% or less. When the amount of retained austenite exceeds 25%, the hardness decreases, and in particular, it is intended to be applied to gear members (for example, gears for power and automobiles (for example, gears for transmissions)) that are easily subjected to heavy loads or impacts. In this case, there is a problem that initial damage such as gear deformation and tooth surface rippling (lip ring) is likely to occur. The amount of retained austenite on the carburized layer surface is more preferably 20% or less.
[0048]
On the other hand, the lower limit of the area ratio of the retained austenite is not particularly limited, and the amount can be reduced by forcibly converting the retained austenite into martensite by, for example, shot peening. In this case, the final retained austenite amount can be reduced to a maximum of about 1% (this is a value that is practically below the measurement limit), for example, in area ratio. In addition, since the martensitic transformation of steel has a large volume expansion, the austenite remaining under the influence of the back stress may not be completely reduced to zero.
[0049]
The amount of retained austenite is the main diffraction peak of the body-centered cubic (ferrite) or body-centered tetragonal (martensite) phase when the X-ray diffraction profile is measured on the surface of the carburized hardened layer by the diffractometer method. Integration area of If ({200} and {211}; since the peak splitting in the tetragonal system, all peak areas belonging to the group are summed) and the integration of the diffraction peaks of the face-centered cubic system, that is, the austenite phase The area is Ia ({200}, {220}, {311}},
{Ia / (Ia + If)} × 100 (%)
It shall be represented by
[0050]
Furthermore, the area ratio of the troostite structure on the surface of the carburized layer is desirably 10% or less. Truthite is an incompletely quenched structure formed in the carburized layer after carburizing and quenching, and has low hardness. Therefore, when the area ratio of the carburized layer surface exceeds 10%, the surface fatigue strength is remarkably deteriorated. Therefore, it is desirable to restrict the area ratio of the above-mentioned troostite to 10% or less particularly when the surface fatigue strength is important.
[0051]
The depth at which grain boundary oxidation from the member surface occurs is preferably within 3 μm. Oxygen that penetrates into the member (steel) from the atmosphere during carburizing is often dominated by grain boundary diffusion, and forms an oxide phase at the grain boundary. When the grain boundary oxide phase is formed, the grain boundary strength is lowered, and the carburized layer may be insufficient in strength, or may be deteriorated in wear resistance due to degranulation. In addition, when the grain boundary oxide phase is generated, some of the surrounding steel additive elements (subcomponents) are taken into the grain boundary oxide phase, resulting in a depletion region of the additive elements around the grain boundary oxide phase, There is a possibility that the hardenability of the charcoal hardened layer is insufficient, and that the hardness and strength are insufficient. Therefore, the grain boundary oxidation depth is suppressed to 3 μm or less by adjusting the composition of the steel material, the atmosphere during carburization (particularly oxygen partial pressure), the carburizing temperature, the time, and the like.
[0052]
The grain boundary oxidation phase can be easily discriminated visually from the non-oxidized region by polishing the cross section of the member. Therefore, the grain boundary oxidation depth can be measured from the cross-sectional optical micrograph.
[0053]
Next, the surface compressive residual stress of the member is desirably 300 to 800 MPa. By causing the compressive residual stress to remain on the surface of the member, when a crack is formed in the surface layer portion, expansion and propagation of the crack are suppressed. Therefore, the strength of the member, particularly dynamic strength (surface fatigue strength, bending fatigue strength, torsional fatigue strength, etc.) can be greatly improved. As described above, when martensite is generated by quenching the carburized layer, a compressive stress field is generated due to volume expansion accompanying transformation, which is convenient for imparting the above surface compressive residual stress. However, when the amount of martensite produced is small, that is, when the retained austenite is large, a sufficient compressive residual stress field cannot be formed. Accordingly, it can be said that reducing the retained austenite (specifically, 25% or less) is advantageous in terms of enhancing the compressive residual stress effect. In addition, absorption of volume expansion at the time of martensite transformation reduces stress because it progresses by plastic deformation of surrounding austenite when the amount of martensite is small, and does not contribute much to increase in surface compressive residual stress. However, when the amount of martensite increases and the retained austenite decreases as described above, the density of dislocations introduced by plastic deformation increases, and slip deformation is constrained, so the surface compressive residual stress increases rapidly. There is also a method of increasing the compressive residual stress by performing surface layer processing such as shot peening after quenching. In the latter case, if the retained austenite is converted to martensite by shot peening, it is more advantageous in improving the compressive residual stress. If the compressive residual stress is less than 300 MPa, the effect of improving the strength by suppressing crack propagation cannot be obtained sufficiently, and imparting a compressive residual stress exceeding 800 MPa increases the amount of martensite or causes strain due to post-processing. One must choose between making the application too large, the former must increase the cooling rate during quenching beyond the limit, and the latter must excessively increase the processing strain. Disappear. In any case, in the light of the problem of the present invention to ensure the dimensional accuracy of the members, the result of falling over at the end is caused.
[0054]
The residual stress is measured when the X-ray diffraction profile is measured by the diffractometer method on the surface of the carburized and quenched layer, and the (211) phase of the body-centered cubic (ferrite) or body-centered tetragonal (martensite) phase is obtained. ) It can be performed based on the relationship between the half width of the peak and the peak center position. An apparatus for performing stress analysis by this method is well-known as a Fastless (registered trademark) stress analyzer (for example, “New Edition X-ray diffraction theory” (Kariti, Agne (1979), pages 431 to 433), Description is omitted.
[0055]
The carburized and quenched member of the present invention can be realized by the manufacturing method of the present invention, and the steel constituting the member non-carburized portion is: subcomponentContent ofC: 0.12-0.22 mass%, Si: 0.4-1.5 mass%, Mn: 0.25-0.45 mass%, Ni: 0.5-1.5 mass%, Cr: 1.3-2.3 mass%, B: 0.001-0.003 mass%, Ti: 0.02-0.06 mass%, Nb: 0.02-0.12 mass%, Al: 0.005 to 0.05 mass%,Mo: 0.05 mass% or less is set, and the balance is set to Fe and inevitable impurities as the main component.;
N≡106 × C (mass%) + 10.8 × Si (mass%) + 19.9 × Mn (mass%) + 16.7 × Ni (mass%) + 8.55 × Cr (mass%) + 45.5 × Mo The component parameter N represented by (mass%) + 28 is 95 or less;
In addition, when cooling at a constant rate from 900 ° C. to room temperature, bainite is not generated in the range where the cooling rate is at least 0.1 ° C./second or less, and in the range where the cooling rate is at least 12 ° C./second or more. To prevent the formation of ferrite;
The subcomponent content is adjusted steel, and a carburized layer is formed on the surface of the member,After quenchingVickers hardness measured on the surface of the carburized layer is 700 Hv or more,After quenchingThe Vickers hardness of the non-carburized portion inside the carburized layer is 250 Hv or more.
[0056]
The carburizing and quenching member of the present invention is easy to process before carburizing and quenching by employing the steel of the above composition, and a sufficient quenching structure can be formed by gas cooling regardless of the member dimensions. A highly accurate member with low strain can be realized. In addition, the fracture strength and surface fatigue strength after carburizing and quenching are excellent, and the manufacturing cost of the entire member can be greatly reduced. Since the critical meaning of each numerical value has already been explained, it will be omitted. Note that the composition of the carburized layer corresponds to a composition in which the carbon content is increased by carburization based on the composition of the non-carburized portion.
[0057]
【Example】
Hereinafter, experimental results performed to confirm the effects of the present invention will be described.
First, steels having the chemical compositions shown in Table 1 were melted in an arc furnace, then formed into round bars having a diameter of 150 mm and a diameter of 32 mm by hot rolling, kept at 925 ° C. for 1 hour, and then air-cooled and normalized. Steel types A1, A2, and A3 are steel types corresponding to the steel composition described in the claims of the present invention, and steel types B and C are steel types corresponding to JIS case-hardened steel SCM420 and SNCM815, respectively (component parameters) The calculated value of N is also shown).
[0058]
[Table 1]
For all steel types, Rockwell hardness was measured at the center of the cross section of the normalized material having a diameter of 32 mm. And the lower limit critical cooling rate (beta) LC of bainite formation and the upper limit critical cooling rate (beta) UC of ferrite formation of each steel material were measured by the method already demonstrated by making the test piece of FIG. 4 separately. Further, a round bar test piece having a diameter of 25 mm and a length of 50 mm and a rotating bending fatigue test piece having the shape shown in FIG. 1 were processed. Moreover, the roller for a pitting test shown in FIG. 2 and its counterpart roller were processed from the normalized material having a diameter of 32 mm and a diameter of 150 mm. On the other hand, the normalizing material having a diameter of 150 mm was cut by a cross section orthogonal to the axis, and a gear having the shape shown in FIG. 3 was processed by forging and cutting.
[0060]
Among these members, for the round bar test pieces having a diameter of 25 mm, the steel types A1, A2, A3, B, and C were subjected to low-pressure carburizing and gas quenching under the conditions shown in Table 2 (invention example A and comparative example, respectively). B, referred to as Comparative Example C). And about a gear, a rotation bending fatigue test piece, a pitting test roller, and a counter roller, steel types A1, A2, A3, and steel type C perform low-pressure carburization and gas quenching on the conditions shown in Table 2 (invention example A1, A2). , A3, Comparative Example C), and Steel type B were subjected to gas carburization and oil quenching under the conditions shown in Table 3.
[0061]
[Table 2]
[Table 3]
Further, the axial cross section of the test piece was mirror-polished and the surface layer portion was subjected to XMA (X-ray Micro Analysis) analysis to examine the carbon concentration of the surface layer portion after carburizing at a position of 50 μm from the surface (result) Are shown in Table 1). On the other hand, the hardness distribution of the cross section of the round bar test piece having a diameter of 25 mm was examined using a Vickers hardness tester. The surface hardness of the carburized quenching material was measured at a position 0.02 mm from the surface. Furthermore, the area ratio of troostite was measured by image analysis of a scanning electron micrograph at the same position. Further, the amount of retained austenite and compressive residual stress in the surface layer portion were measured by the method described above. These results are shown in Table 4.
[0064]
[Table 4]
In Table 4, Invention Examples A1 to A3 are 95HRBThe following normalization hardness is shown, and the hardness of the center portion of the carburized quenching material is 350 Hv or more. The structure of the surface layer (carburized layer) is martensite, and the structure of the central part (non-carburized part) is bainite or bainite + martensite (however, bainite is 50% or more), which is remarkable incomplete firing. There was no institution. On the other hand, although the normalization hardness is low in Comparative Example B, the surface layer hardness and the center part hardness of the carburized quenching material are lower than those of the invention examples. In Comparative Example C, the surface hardness and the center hardness of the carburized and quenched material are substantially the same as those of Invention Example A, but the normalizing hardness is extremely high.
[0066]
In the rotating bending fatigue test, the Ono type rotating bending fatigue tester was used to determine the fatigue strength based on the number of repetitions of 10 million times. Also, in the pitting test, the test roller and the mating roller are rotated and contacted in the lubricant while controlling the load and sliding speed, and the pitting damage generated on the surface of the test roller is detected by vibration, and the number of repetitions is 10 million times. The surface pressure strength with respect to was determined. The relative sliding speed was 800 mm / s. Table 4 shows the results of these strength tests (however, each strength is expressed as a relative value based on Comparative Example B). As can be seen from the results, the invention examples have characteristics that greatly exceed the comparative examples in the rotational bending fatigue strength and the anti-pitting strength.
[0067]
Next, regarding the gear, the accuracy after carburizing and quenching was evaluated by the following method. First, regarding the gear accuracy, the gears were set in a dedicated precision gear accuracy measuring machine, and the error amount in the pressure angle direction and the error amount in the torsion angle direction of the gear were measured on each of the left and right tooth surfaces. Further, the height of the tooth gap was measured over the entire circumference, and a value obtained by subtracting the minimum value from the maximum value was calculated as the tooth gap runout. On the other hand, the dimensional accuracy of the gears was measured as follows. That is, a ball is put into two tooth spaces that are facing each other of the gear, and the outer peripheral dimensions thereof are set to a dedicated O.D. B. It measured with D (Over Ball Diaphragm) measuring device. O. B. As shown in FIG. 3, the measurement position of D was two directions (X, Y) at right angles in the circumferential direction, and three positions (A, B, C) at the upper, middle, and lower sides in the tooth width direction. In the table, “OBD ellipse” means O.D. B. It is the absolute value of the difference between D and “OBD taper” is the upper O.D. B. D and lower O.D. B. D difference. The results are shown in Table 5.
[0068]
[Table 5]
That is, it can be seen that the gear configured as the carburized and quenched member according to the present invention has very good characteristics over all the gear accuracy and dimensional accuracy because the heat treatment strain is extremely small.
[Brief description of the drawings]
FIG. 1 shows a rotating bending fatigue test piece.
FIG. 2 is a view showing a pitting test roller and a mating roller.
FIG. 3 is a view showing a heat treatment strain evaluation gear.
FIG. 4 is a diagram showing the shape of a sample for continuous cooling transformation diagram measurement.
Claims (16)
前記ガス焼入れを行なう際のガス冷却雰囲気は、前記部材を構成する鋼の熱伝導率をλ（単位：ｋｃａｌ／ｍｈ℃）、該ガス冷却雰囲気における前記部材の表面熱伝達係数をα（単位：ｋｃａｌ／ｍｈ２℃）として、Ｈ≡０．５×（α／λ）にて定義される焼入急冷度Ｈが０．０１〜０．０８となる雰囲気が使用されることを特徴とする浸炭焼入れ部材の製造方法。 The content of subcomponents is C: 0.12-0.22 mass%, Si: 0.4-1.5 mass%, Mn: 0.25-0.45 mass%, Ni: 0.5-1 0.5% by mass, Cr: 1.3-2.3% by mass, B: 0.001-0.003% by mass, Ti: 0.02-0.06% by mass, Nb: 0.02-0.12 Mass%, Al: 0.005 to 0.05 mass%, Mo: 0.05 mass% or less, the balance is set to Fe and inevitable impurities, and the constant rate from 900 ° C. to room temperature So that bainite is not generated when the cooling rate is at least 0.1 ° C./second or less, and ferrite is not generated when the cooling rate is at least 12 ° C./second or more. The member is made of steel in which the content of each subcomponent in the non-carburized state is adjusted. And carburizing the member in a reduced pressure atmosphere of 1 to 30 HPa to form a carburized layer having a surface carbon concentration of 0.6 to 1.5% by mass on the member, and after the carburizing treatment, the member Is quenched by gas cooling using an inert gas or nitrogen gas so that the quenching quenching degree H becomes 0.01 to 0.08,
In the gas cooling atmosphere when performing the gas quenching, the thermal conductivity of the steel constituting the member is λ (unit: kcal / mh ° C.), and the surface heat transfer coefficient of the member in the gas cooling atmosphere is α (unit: kcal / mh2 ° C.), carburizing and quenching characterized in that an atmosphere in which quenching quenching degree H defined by H≡0.5 × (α / λ) is 0.01 to 0.08 is used. Manufacturing method of member.
Ｎ≡１０６×Ｃ（質量％）＋１０．８×Ｓｉ（質量％）＋１９．９×Ｍｎ（質量％）＋１６．７×Ｎｉ（質量％）＋８．５５×Ｃｒ（質量％）＋４５．５×Ｍｏ（質量％）＋２８により表される成分パラメータＮが９５以下となり； N≡106 × C (mass%) + 10.8 × Si (mass%) + 19.9 × Mn (mass%) + 16.7 × Ni (mass%) + 8.55 × Cr (mass%) + 45.5 × Mo The component parameter N represented by (mass%) + 28 is 95 or less;
また、９００℃から室温まで一定速度にて冷却したとき、冷却速度が少なくとも０．１ Further, when cooling at a constant rate from 900 ° C. to room temperature, the cooling rate is at least 0.1. ℃／秒以下の範囲においてはベイナイトが生成せず、かつ、冷却速度が少なくとも１２℃／秒以上の範囲においてはフェライトが生成しなくなるように；Bainite is not generated in the range of ° C / second or less, and ferrite is not generated in the range of the cooling rate of at least 12 ° C / second or higher;
前記副成分含有量が調整された鋼であり、前記部材表面に浸炭層が形成されるとともに、焼入れ後の該浸炭層の表面にて測定したビッカース硬度が７００Ｈｖ以上であり、また、焼入れ後の浸炭層内側の非浸炭部のビッカース硬度が２５０Ｈｖ以上であることを特徴とする浸炭焼入れ部材。 It is a steel with an adjusted content of subcomponents, and a carburized layer is formed on the surface of the member, and a Vickers hardness measured on the surface of the carburized layer after quenching is 700 Hv or more, and after quenching A carburized and quenched member, wherein a non-carburized portion inside the carburized layer has a Vickers hardness of 250 Hv or more.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001391133A JP4050512B2 (en) | 2001-12-25 | 2001-12-25 | Manufacturing method of carburizing and quenching member and carburizing and quenching member |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001391133A JP4050512B2 (en) | 2001-12-25 | 2001-12-25 | Manufacturing method of carburizing and quenching member and carburizing and quenching member |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2003193128A JP2003193128A (en) | 2003-07-09 |
| JP4050512B2 true JP4050512B2 (en) | 2008-02-20 |
Family
ID=27598814
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2001391133A Expired - Fee Related JP4050512B2 (en) | 2001-12-25 | 2001-12-25 | Manufacturing method of carburizing and quenching member and carburizing and quenching member |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP4050512B2 (en) |
Cited By (2)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US10156006B2 (en) | 2009-08-07 | 2018-12-18 | Swagelok Company | Low temperature carburization under soft vacuum |
| US10246766B2 (en) | 2012-01-20 | 2019-04-02 | Swagelok Company | Concurrent flow of activating gas in low temperature carburization |
Families Citing this family (9)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP3981068B2 (en) * | 2003-12-25 | 2007-09-26 | 株式会社神戸製鋼所 | Method for producing galvannealed steel sheet |
| JP4964001B2 (en) * | 2007-03-29 | 2012-06-27 | 株式会社神戸製鋼所 | Gears with excellent low cycle fatigue strength |
| JP5361293B2 (en) * | 2008-09-03 | 2013-12-04 | 山陽特殊製鋼株式会社 | Carburized parts with excellent static strength |
| CN103290419B (en) * | 2013-06-25 | 2015-04-15 | 清华大学 | Gear treatment method |
| JP6308382B2 (en) * | 2013-11-21 | 2018-04-11 | 大同特殊鋼株式会社 | Carburized parts |
| JP2015134945A (en) * | 2014-01-16 | 2015-07-27 | 山陽特殊製鋼株式会社 | Carburizing steel |
| JP7510045B2 (en) * | 2020-03-27 | 2024-07-03 | 日本製鉄株式会社 | Steel for carburizing, carburizing steel parts, and method for manufacturing carburizing steel parts |
| CN112695269B (en) * | 2020-11-30 | 2022-09-27 | 山西平阳重工机械有限责任公司 | Heat treatment process of 18Cr2Ni4WA workpiece |
| CN113046524B (en) * | 2021-03-05 | 2024-02-27 | 苏州奥轩精密科技有限公司 | Traction ring heat treatment process |
-
2001
- 2001-12-25 JP JP2001391133A patent/JP4050512B2/en not_active Expired - Fee Related
Cited By (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US10156006B2 (en) | 2009-08-07 | 2018-12-18 | Swagelok Company | Low temperature carburization under soft vacuum |
| US10934611B2 (en) | 2009-08-07 | 2021-03-02 | Swagelok Company | Low temperature carburization under soft vacuum |
| US10246766B2 (en) | 2012-01-20 | 2019-04-02 | Swagelok Company | Concurrent flow of activating gas in low temperature carburization |
| US11035032B2 (en) | 2012-01-20 | 2021-06-15 | Swagelok Company | Concurrent flow of activating gas in low temperature carburization |
Also Published As
| Publication number | Publication date |
|---|---|
| JP2003193128A (en) | 2003-07-09 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| JP5958652B2 (en) | Soft nitrided induction hardened steel parts with excellent surface fatigue strength | |
| JP5994924B2 (en) | Shaped material of induction-hardened parts and method for manufacturing the same | |
| JP4956146B2 (en) | Case-hardened steel excellent in forgeability and prevention of grain coarsening, its manufacturing method, and carburized parts | |
| WO2011122651A1 (en) | Carburized steel member and method for producing same | |
| CN107406959B (en) | Nitrided component and soft nitrided component having excellent wear resistance and pitting resistance, and nitriding method and soft nitriding method | |
| JP5872863B2 (en) | Gear having excellent pitting resistance and method for producing the same | |
| JP5477111B2 (en) | Nitriding induction hardening steel and nitriding induction hardening parts | |
| WO2012105405A1 (en) | Steel for nitriding and nitrided component | |
| JP4354277B2 (en) | Method for manufacturing carburized and quenched members | |
| WO2019198539A1 (en) | Machine component and method for producing same | |
| JP4050512B2 (en) | Manufacturing method of carburizing and quenching member and carburizing and quenching member | |
| JP6314648B2 (en) | Surface hardened component and method for manufacturing surface hardened component | |
| JP2017171951A (en) | Steel component and production method thereof | |
| WO2015146837A1 (en) | Case-hardened steel having excellent cold forgeability and capable of suppressing abnormal grain growth during carburizing treatment | |
| JP2017133052A (en) | Case hardened steel excellent in coarse particle prevention property, fatigue property and machinability during carburization and manufacturing method therefor | |
| JP5336972B2 (en) | Nitriding steel and nitride parts | |
| JP7013833B2 (en) | Carburized parts | |
| JP6447064B2 (en) | Steel parts | |
| JP7063070B2 (en) | Carburized parts | |
| JP7006052B2 (en) | Steel material for soaking treatment | |
| WO2017146057A1 (en) | Cement steel component and steel material having excellent stability of rolling fatigue life, and method for manufacturing same | |
| JP2016188422A (en) | Carburized component | |
| JP7510045B2 (en) | Steel for carburizing, carburizing steel parts, and method for manufacturing carburizing steel parts | |
| JP2017150066A (en) | Steel material and carburized steel component excellent in stability of rolling motion fatigue life and manufacturing method therefor | |
| JP7273324B2 (en) | Nitrided part blanks and nitrided parts |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20040510 |
|
| A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20070808 |
|
| A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20070827 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20071024 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20071115 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20071129 |
|
| R150 | Certificate of patent or registration of utility model |
Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20101207 Year of fee payment: 3 |
|
| S111 | Request for change of ownership or part of ownership |
Free format text: JAPANESE INTERMEDIATE CODE: R313117 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20101207 Year of fee payment: 3 |
|
| R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20111207 Year of fee payment: 4 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20111207 Year of fee payment: 4 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20121207 Year of fee payment: 5 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20131207 Year of fee payment: 6 |
|
| LAPS | Cancellation because of no payment of annual fees |