JPWO2002044435A1

JPWO2002044435A1 - Carburizing steel and carburizing gear

Info

Publication number: JPWO2002044435A1
Application number: JP2002546781A
Authority: JP
Inventors: 野村　一衛; 加藤　智也; 住田　庸; 伊藤　幸夫; 大西　昌澄; 相原　秀雄; 三林　雅彦; 江里口　正
Original assignee: Toyota Motor Corp; Aichi Steel Corp
Current assignee: Toyota Motor Corp; Aichi Steel Corp
Priority date: 2000-12-01
Filing date: 2001-11-30
Publication date: 2004-04-02
Anticipated expiration: 2021-11-30
Also published as: JP4136656B2; WO2002044435A1

Abstract

本発明の目的は、冷間鍛造成形性に優れる浸炭用鋼及び低サイクル疲労強度に優れる浸炭歯車を提供することである。本発明の浸炭用鋼は、重量％で、Ｃ：０．１０〜０．３０％，Ｓｉ：０．５０％以下，Ｍｎ：０．５０〜１．５０％，Ｐ：０．０３０％以下，Ｓ：０．０３０％以下，Ｃｒ：０．８５〜２．００％，Ｍｏ：０．３５％以下，Ｂ：０．００１０〜０．００５０％，Ａｌ：０．１１〜０．３０％，Ｎ：０．００８０〜０．０２５０％，Ｎｂ：０．０１〜０．１０％，Ｔｉ：０．０１〜０．１０％を含有し、残部はＦｅ及び不可避不純物からなり、Ｈｄ＝８３Ｃ重量％＋５．５Ｍｎ重量％＋４．０Ｃｒ重量％＋１０．５Ｍｏ重量％＋１２の関係式において、Ｈｄ≧６０√Ｃ重量％＋１２．５であり、有効Ａｌ重量％＝Ａｌ重量％−２（Ｎ重量％−０．３０Ｔｉ重量％）≧０．１である。An object of the present invention is to provide a carburizing steel excellent in cold forgeability and a carburized gear excellent in low cycle fatigue strength. The carburizing steel of the present invention is, by weight%, C: 0.10 to 0.30%, Si: 0.50% or less, Mn: 0.50 to 1.50%, P: 0.030% or less, S: 0.030% or less, Cr: 0.85 to 2.00%, Mo: 0.35% or less, B: 0.0010 to 0.0050%, Al: 0.11 to 0.30%, N : 0.0080 to 0.0250%, Nb: 0.01 to 0.10%, Ti: 0.01 to 0.10%, the balance being Fe and unavoidable impurities, Hd = 83C weight% + 5 In a relational expression of 0.5 Mn wt% + 4.0 Cr wt% + 10.5 Mo wt% + 12, Hd ≧ 60 ° C. wt% + 12.5, and effective Al wt% = Al wt% −2 (N wt% −0. 30Ti wt%) ≧ 0.1.

Description

技術分野
本発明は、浸炭用鋼及びそれを用いた浸炭歯車に関し、更に詳しくは、冷間鍛造成形性に優れる浸炭用鋼及び低サイクル疲労強度に優れる浸炭歯車に関する。
背景技術
従来、自動車、産業用機械等に使用される歯車は、ＪＩＳ　ＳＭｎＣ，ＳＣｒ，ＳＣＭ，ＳＮＣＭ等の機械構造用合金鋼を素材として、これらを熱間鍛造、温間鍛造又は冷間鍛造後、機械加工を施して成形した後、耐磨耗性や疲労強度を向上させるため表面硬化処理（浸炭焼入、高周波焼入、軟窒化処理等）を施してから供されている。
特公平７−１００８４０号公報にはＭｏ等の合金元素を多量に添加した鋼を用いた歯車が示され、これにより低サイクル疲労強度（スポーリング破壊等の歯面疲労強度や衝撃的な歯元疲労強度等）が改善されることが知られている。しかし、高価な合金元素の添加に加えて、冷間鍛造成形性や機械加工性を悪化させるため、コストを大幅に上昇させる等の欠点があった。そのため、特開平１０−１５２７４６号公報及び特開平１１−７１６５４号公報に示されるように、素材に添加する合金成分を削減し、これによる焼入性の低下を補うため少量のＢを添加し、加工性と強度を両立させる浸炭用鋼が提案されている。
しかしながら、この２件の公報に記載の発明は、歯車の特に浸炭層のトルースタイト生成防止と浸炭時のオーステナイト結晶粒の安定性に関する検討が十分でなく、優れた低サイクル疲労強度が得られない。
尚、ここでトルースタイトとは、通常ガス浸炭等を行った場合、歯車表面に観察される不完全焼入層（浸炭ガス雰囲気の酸素が歯車表面から拡散し、素材に含まれるＳｉ，Ｍｎ，Ｃｒ等の合金元素と酸化物を形成するため、その周辺部のＳｉ，Ｍｎ，Ｃｒ等の固溶合金元素が欠乏するため、焼入性が低下して生じる組織）とは異なり、浸炭焼入冷却時に主に浸炭層のオーステナイト粒界に沿って析出する微細パーライトを指し、これが多量に存在すると低サイクル疲労強度が低下する大きな要因となる。
本発明は、上記実情に鑑みてなされたものであり、冷間鍛造成形性に優れる浸炭用鋼及び低サイクル疲労強度に優れる浸炭歯車を提供することを目的とする。
発明の開示
本発明者らは、冷間鍛造成形性に優れる浸炭用鋼及び低サイクル疲労強度に優れる浸炭歯車について検討した結果、本発明を完成するに至った。
即ち、本発明の浸炭用鋼は、Ｃ：０．１０〜０．３０重量％，Ｓｉ：０．５０重量％以下，Ｍｎ：０．５０〜１．５０重量％，Ｐ：０．０３０重量％以下，Ｓ：０．０３０重量％以下，Ｃｒ：０．８５〜２．００重量％，Ｍｏ：０．３５重量％以下，Ｂ：０．００１０〜０．００５０重量％，Ａｌ：０．１１〜０．３０重量％，Ｎ：０．００８０〜０．０２５０重量％，Ｎｂ：０．０１〜０．１０重量％，Ｔｉ：０．０１〜０．１０重量％を含有し、残部はＦｅ及び不可避不純物からなり、
Ｈｄ＝８３Ｃ重量％＋５．５Ｍｎ重量％＋４．０Ｃｒ重量％＋１０．５Ｍｏ重量％＋１２の関係式において、Ｈｄ≧６０√Ｃ重量％＋１２．５であり、且つ有効Ａｌ重量％＝Ａｌ重量％−２（Ｎ重量％−０．３０Ｔｉ重量％）≧０．１であることを特徴とする。
上記Ｃは、浸炭焼入後の肌焼深さ及び内部硬さを向上させ、低サイクル疲労強度の向上に大きく影響する。しかし、Ｃの含有量が０．１０重量％未満では、肌焼深さを得るのに長い浸炭時間を要するためコスト高となり、また、内部強度も低下して歯車における低サイクル疲労強度が大きく低下する。一方、０．３０重量％を超えると、浸炭焼入後の内部硬さが高くなり、歯車の歯の靱性が著しく低下するため、低サイクル疲労強度が低下する。加えて、歯車の製造工程である冷間加工性や被削性が悪くなり、金型や工具等の寿命を劣化させるため、著しいコスト高となる。尚、Ｃの含有量の下限は、好ましくは０．１３重量％、より好ましくは０．１９重量％であり、Ｃの含有量の上限は、好ましくは０．２７重量％、より好ましくは０．２４重量％である。
上記Ｓｉは脱酸剤として添加される元素であるが、フェライトに固溶して強化する元素であるため、特に冷間加工性に影響する。Ｓｉの含有量が０．５０重量％を超えると、浸炭性を阻害するおそれがある。Ｓｉの含有量は、好ましくは０．３５重量％以下、より好ましくは０．１５重量％以下である。尚、下限は通常０．０３重量％である。
上記Ｍｎは脱酸剤として添加される元素であるが、素材の焼入性を向上させるため、浸炭焼入後の肌焼深さ及び内部硬さを向上させ、低サイクル疲労強度向上に大きく影響する。しかし、Ｍｎの含有量が０．５０重量％未満では、浸炭拡散層のトルースタイト生成が顕著になり、内部硬さも低下して歯車における低サイクル疲労強度が大きく低下する。一方、１．５０重量％を超えると、軟化焼鈍時のパーライトの変態温度が下がるためにパーライト部の硬さがＨｖ３００を超えるため、歯車の冷間加工性や被削性が悪くなり、金型や工具等の寿命を劣化させるため、著しいコスト高となる。尚、Ｍｎの含有量の下限は、好ましくは０．８重量％、より好ましくは１．０重量％であり、Ｍｎの含有量の上限は、好ましくは１．５重量％、より好ましくは１．４重量％である。
上記Ｐはフェライトに固溶して強化する元素であり、特に冷間加工性を劣化させる元素である。Ｐの含有量が０．５重量％を超えると浸炭時のオーステナイト粒界に偏析し、浸炭層の粒界強度が低下する。Ｐの含有量は、好ましくは０．０１５重量％以下、より好ましくは０．０１２重量％以下である。尚、下限は、通常０．００２重量％である。
上記Ｓは多量に含有すると浸炭時のオーステナイト粒界に偏析し、浸炭層の粒界強度を低下させる。また、Ｍｎとの化合物であるＭｎＳを形成するため、冷間加工性を低下させる元素でもあるが、一方歯車の被削性を向上させる元素でもある。Ｓの含有量の下限は、好ましくは０．００５重量％、より好ましくは０．００８重量％であり、Ｓの含有量の上限は、好ましくは０．０２０重量％、より好ましくは０．０１５重量％である。
上記Ｃｒは素材の焼入性を向上させるため、浸炭焼入後の肌焼深さ及び内部硬さを向上させ、低サイクル疲労強度の向上に大きく影響する。特に、Ｂが添加される鋼においては、浸炭拡散層に低サイクル疲労強度を低下させるトルースタイトが生成しやすくなるが、その生成を防止するために著しい効果がある。Ｃｒの含有量が０．８５重量％未満では浸炭層のトルースタイト生成が顕著となり、一方、２．００重量％を超えると浸炭時のオーステナイト粒界にＣｒ炭化物が析出し、浸炭層の粒界強度を低下させることとなり、好ましくない。Ｃｒの含有量の下限は、好ましくは０．９重量％、より好ましくは１．０重量％であり、Ｃｒの含有量の上限は、好ましくは１．５重量％、より好ましくは１．３重量％である。
上記Ｍｏは素材の焼入性を向上させ、浸炭焼入後の肌焼深さ及び内部硬さを向上させ、更に低サイクル疲労強度の向上に効果がある。しかし、Ｍｏの含有量が０．３５重量％を超えると、歯車の製造中に硬さが著しく上がり、冷間加工性や被削性が悪くなり、金型や工具等の寿命を短くし、コスト高となってしまう。また、Ｂが添加される鋼においては、Ｍｏを多く含むと浸炭層にトルースタイトが生成しやすくなる。Ｍｏの含有量は、好ましくは０．２５重量％以下、より好ましくは０．２０重量％以下である。尚、下限は通常０．００５重量％である。
上記Ｂは浸炭層のオーステナイト粒界に析出して浸炭層の粒界強度を向上させ、冷間鍛造時の素材硬さを上昇させずに浸炭焼入後の内部の焼入性を向上させるために重要な元素である。しかし、浸炭層の焼入性向上効果が小さいために含有量が０．００５０重量％を超えると、浸炭層にトルースタイトを生成しやすくなり、更には焼入性向上効果が飽和するだけでなく、熱間鍛造成形性又は冷間鍛造成形性が悪化し、０．００１０重量％未満では浸炭焼入時の内部の焼入性が低下し好ましくない。Ｂの含有量は、好ましくは０．００１０〜０．００３５重量％、より好ましくは０．００１５〜０．００３０重量％である。
上記Ａｌは脱酸剤として添加される元素であるが、鋼中のＮと反応してＡｌＮを形成し、浸炭加熱時のオーステナイト結晶粒の粗大化を防止する作用がある。また、有効Ａｌ重量％＝Ａｌ重量％−２（Ｎ重量％−０．３０Ｔｉ重量％）≧０．１において、Ａｌの含有量が０．１１重量％未満では、浸炭層にトルースタイトが生成しやすくなり、且つ、冷間加工率の高い歯車の場合における浸炭時のオーステナイト結晶粒の異常粒の成長が見られ、Ａｌの含有量が０．３重量％を超えると浸炭時に歯車表層の不完全焼入層が深くなり好ましくない。また、有効Ａｌ重量％が０．１重量％未満では、焼入性が低下し、浸炭層のトルースタイトが顕著になる。Ａｌの含有量は、好ましくは０．１１〜０．２０重量％、更に好ましくは０．１１〜０．１５重量％である。
また、形成されるＡｌＮの含有量は、好ましくは２００〜６００ｐｐｍ、より好ましくは２５０〜５５０ｐｐｍ、更に好ましくは３００〜５００ｐｐｍである。
上記ＮはＡｌと反応してＡｌＮを析出するが、Ｎの含有量が０．００８０重量％未満では、冷間加工された歯車において、析出するＡｌＮ量が低下して浸炭時のオーステナイト結晶粒が粗大化する。一方、０．０２５０重量％を超えると固溶Ｂ量が減少し、焼入性が低下してしまう。Ｎの含有量の下限は、好ましくは０．００８重量％、より好ましくは０．０１２重量％、Ｎの含有量の上限は、好ましくは０．０２５重量％、より好ましくは０．０２０重量％である。
上記ＮｂはＮ又はＣと反応してＮｂ（Ｃ，Ｎ）を生成し、浸炭加熱時のオーステナイト結晶粒の粗大化を防止する効果がある。Ｎｂの含有量が０．０１重量％未満では、上記効果が発現せず、また、０．１０重量％を超えると浸炭加熱時のオーステナイト結晶粒の粗大化を防止する効果が飽和し、コスト高となり好ましくない。Ｎｂの含有量は、好ましくは０．０１〜０．０７重量％、より好ましくは０．０１〜０．０５重量％である。
上記ＴｉはＮと反応してＴｉＮを生成し、浸炭加熱時のオーステナイト結晶粒の粗大化を防止する効果がある。Ｔｉの含有量が０．０１重量％未満では、上記効果が発現せず、また、０．１０重量％を超えると浸炭加熱時のオーステナイト結晶粒の粗大化を防止する効果が飽和し、コスト高となり好ましくない。Ｔｉの含有量は、好ましくは０．０１〜０．０７重量％、より好ましくは０．０１〜０．０５５重量％である。
上記式Ｈｄ＝８３Ｃ重量％＋５．５Ｍｎ重量％＋４．０Ｃｒ重量％＋１０．５Ｍｏ重量％＋１２は、上記組成を有する浸炭用鋼の浸炭焼入後における内部硬さの大小を表わすパラメーターを求める式であり、この値が大きいほど浸炭焼入後の内部硬さが高くなる。
このようにして求められたＨｄが上記式Ｈｄ≧６０√Ｃ重量％＋１２．５を満たした場合、浸炭焼入焼戻の後にマルテンサイト率を９０％以上とすることができる。
また、上記のような浸炭用鋼を浸炭焼入する前に所定形状に加工する際において適切な軟化焼鈍を施した素材状態では、硬さがＨｖ１７０以下（好ましくはＨｖ１６５以下、より好ましくはＨｖ１６０以下）のフェライト＋パーライト組織とすることができ、且つ荷重１０ｇのマイクロビッカース硬さにおいて、パーライト組織の硬さをＨｖ３００以下（好ましくはＨｖ２９５以下、より好ましくはＨｖ２９０以下）とすることができる。これによって浸炭用鋼の冷間鍛造性が向上する。
本発明の浸炭歯車は、上記記載の浸炭用鋼を用いて製造され、硬さＨｖ５１３である肌焼深さが表面から０．５ｍｍ以上であり、且つ肌焼深さ内のトルースタイト面積率が５％以下であることを特徴とする。
上記浸炭歯車の加工方法は特に限定されないが、通常、熱間鍛造、温間鍛造及び冷間鍛造のいずれかが行われた後、機械加工され、浸炭焼入がなされる。その場合の表面層の硬さは、通常ビッカース硬さで表され、硬さＨｖ５１３である肌焼深さは、好ましくは表面から０．６ｍｍ以上、より好ましくは０．６５ｍｍ以上、更に好ましくは０．７ｍｍ以上である。但し、上限は、通常１．５ｍｍ、より好ましくは１．２ｍｍである。上記肌焼深さが０．５ｍｍ未満であれば歯車内部を起点として破壊することとなり好ましくない。一方、上記肌焼深さが０．５ｍｍ以上であれば疲労強度に著しく優れたものとなる。尚、歯車の硬さの測定方法は実施例で示す。
上記歯車の低サイクル疲労による破壊には、歯面への高い接触応力による歯面内部の塑性変形が原因となる歯面疲労剥離と、衝撃的な歯元への高い曲げ応力により歯元内部が塑性し歯元表面応力が増加するために起こる歯元衝撃疲労破壊とがある（図１参照）。一般に、浸炭焼入した歯車は表層のＣ濃度０．８％から歯内部（未浸炭部）に向かって緩やかなＣ濃度分布をもつが、低サイクル疲労強度が要求される歯車では、この歯内部（未浸炭部）の強度・靱性が重要となるため、浸炭層のトルースタイト生成の抑制と靱性に起因するオーステナイト結晶粒の安定性に加え、内部強度に起因する歯内部の焼入組織（マルテンサイト率）が重要となる。
上記「トルースタイトの面積分率」は、浸炭焼入がなされた後の表面から深さ０．５ｍｍまでの領域における平均面積分率をいい、好ましくは４．７％以下、より好ましくは４．５％以下である。５％を超えると疲労強度が低下し、歯車の歯面（歯車の噛み合わせで互いの歯どうしが接触する部分）における剥離の発生の恐れがある。
上記浸炭歯車のピッチ部の内部硬さは、Ｈｖ３５０以上とすることができ、好ましくはＨｖ３７０以上、より好ましくはＨｖ３９０以上である。Ｈｖ３５０未満では歯面疲労強度が低下し好ましくない。
また、上記浸炭歯車の歯元部の内部硬さは、Ｈｖ３００以上とすることができ、好ましくはＨｖ３３０以上、より好ましくはＨｖ３６０以上である。Ｈｖ３００未満では歯元衝撃疲労強度が低下し好ましくない。尚、上記「歯元部」とは、歯のピッチ円部から歯底に向かうＲ部のことをいう。
上記ピッチ部及び歯元部の内部硬さの好ましい組み合わせとしては、ピッチ部がＨｖ３５０以上、且つ歯元部がＨｖ３００以上、より好ましくは、ピッチ部がＨｖ３７０以上、且つ歯元部がＨｖ３３０以上、更に好ましくは、ピッチ部がＨｖ３９０以上、且つ歯元部がＨｖ３６０以上である。
更に、上記ピッチ部の内部硬さ、歯元部の内部硬さ及びトルースタイト面積率の好ましい組み合わせとしては、（１）ピッチ部の硬さがＨｖ３５０以上、且つ歯元部の硬さがＨｖ３００以上、且つトルースタイト面積率が５％以下、（２）より好ましくは、ピッチ部の硬さがＨｖ３７０以上、且つ歯元部の硬さがＨｖ３３０以上、且つトルースタイト面積率が４．８％以下、（３）更に好ましくは、ピッチ部の硬さがＨｖ３９０以上、且つ歯元部の硬さがＨｖ３６０以上、且つトルースタイト面積率が４．７％以下、（４）特に好ましくは、ピッチ部の硬さがＨｖ４００以上、且つ歯元部の硬さがＨｖ３９０以上、且つトルースタイト面積率が４．５％以下である。
上記浸炭用鋼を用いて浸炭歯車を製造する場合には、通常の浸炭焼入焼戻を行えば、上記の組織及び性質を有するものを得ることができる。特に組織は、歯車の質量及び浸炭焼入媒体の影響を受けることになるが、歯車質量が丸棒相当で外径が好ましくは４０ｍｍ以下、より好ましくは３０ｍｍ以下、更に好ましくは２５ｍｍ以下である。また、焼入媒体は２００℃以下の油焼入であることが好ましい。但し、浸炭雰囲気は、ガス浸炭、浸炭浸窒及び真空浸炭等、いずれの雰囲気で行ってもよい。
本発明の浸炭用鋼によれば、特定の元素組成を有し、Ｈｄ＝８３Ｃ重量％＋５．５Ｍｎ重量％＋４．０Ｃｒ重量％＋１０．５Ｍｏ重量％＋１２の関係式において、Ｈｄ≧６０√Ｃ重量％＋１２．５とすることで、マルテンサイト率９０％以上の浸炭焼入組織を得ることができ、更にはオーステナイト結晶粒の異常成長を防止することができる。また、浸炭焼入前の所定形状への加工の際には適切な軟化焼鈍を施すことにより冷間鍛造成形性を改善し、容易に加工することができる。従って、この浸炭用鋼を用いて自動車や産業機械等に使用される歯車やシャフト等を容易に製造することができる。また、本発明の浸炭歯車によれば、冷間鍛造歯車の効果である歯元部の応力集中の緩和や歯形形状に沿った鍛流線が得られるため、より高強度であり、更に高トルク領域の抵抗性である低サイクル疲労強度（耐歯面疲労強度及び耐歯元疲労強度）に優れる。
発明を実施するための最良の形態
以下に本発明について実施例を挙げて具体的に説明する。
１．浸炭用鋼の製造及び評価
（１）浸炭用鋼の組成から求めた硬さ及び有効Ａｌ重量％について
表１〜３に示す化学組成の鋼Ａ〜Ｓ（Ａ〜Ｈを本発明鋼、Ｉ〜Ｗを比較鋼とした。比較鋼のうちＱ１〜３３鋼及びＲ１〜２０鋼は、それぞれ特開平１０−１５２７４６号公報及び特開平１１−７１６５４号公報に相当するものである。また、Ｓ鋼は低サイクル疲労強度が要求される歯車に多く使用されているＪＩＳ規格のＳＮＣＭ４２０に相当する。）のそれぞれ３０ｋｇを真空溶解炉によって溶製した。
表４及び５に、各成分量から計算されたＨｄ_１＝８３Ｃ重量％＋５．５Ｍｎ重量％＋４．０Ｃｒ重量％＋１０．５Ｍｏ重量％＋１２と、Ｈｄ_２＝６０√Ｃ重量％＋１２．５及びこれらの差Ｈｄ_１−Ｈｄ_２を示した。また、表６及び７に有効Ａｌ重量％の計算結果を示した。尚、表４〜７において、構成元素の含有量が本発明の範囲外のものは数値の横に＊印を添えた。

（２）浸炭用鋼の製造及び評価
Ａ〜Ｐ，Ｑ１，Ｒ１及びＳの鋼塊を１２００℃以上で０．５時間加熱した後、１０００〜１２００℃の温度で熱間鍛造成形を行い直径３０ｍｍの丸棒を製造した。これを９００℃から６００℃まで降温速度７５℃／分で焼鈍して浸炭用鋼とした。
上記で得られた浸炭用鋼のうち、Ａ〜Ｈ，Ｏ，Ｐ及びＳの浸炭用丸棒について、硬さ（素地及びパーライト部の硬さ）及び冷間加工性（７０％据込時の変形抵抗及び限界加工率）を以下の方法で測定した。その結果を表８に示す。
素地硬さは、ビッカース硬さ計（型式；ＡＶＫ−Ｃ２、メーカー名；ＡＫＡＳＨＩ社製）を用いて荷重１０ｋｇで測定した。また、パーライト部の硬さはマイクロビッカース硬さを荷重１０ｇで測定した。
変形抵抗は、直径１０ｍｍ、高さ１５ｍｍの試験片（切欠無）を１００ｔ万能試験装置（型式；ＲＨ−１００、メーカー名；島津社製）を用いてロードセル移動速度１ｍｍ／分で負荷し、７０％据込時の圧縮荷重を測定し、日本塑性加工学会が提案している端面拘束圧縮による変形抵抗測定法を用いて求めた変形抵抗を表８に示した（塑性加工春季講演会予稿集（１９８０）ｐ．５２９〜５３２参照）。
限界加工率は上記試験片を上記装置を用いて端面拘束試験を行い、ロードセル移動速度１ｍｍ／分で負荷し、端部（円周部）に割れが入ったときの据込率を限界加工率とした（塑性と加工Ｎｏ．２４１、（１９８１）ｐ．１３９〜１４４、日本塑性加工学会発行　参照）。

２．浸炭歯車の製造及び評価
上記で得られた浸炭用鋼のうち、Ａ〜Ｎ，Ｑ１，Ｒ１及びＳ〜Ｗについて、浸炭用丸棒を冷間鍛造にてモジュールが４．８、歯数が１０、ピッチ円径が４８．８ｍｍである傘歯車を、図２に示すヒートパターンで浸炭焼入焼戻を行い、更に内径等の研削仕上げを行い製造した。この傘歯車を図３に示すように油圧疲労試験装置にセットし、以下の要領で傘歯車の疲労試験を行った。試験歯車に実相手歯車相当の曲率を持った治具を作製し、完全片振りの疲労試験を行った。また、アコースティック・エミッションを用いて亀裂発生時期を寿命とし、磁粉探傷により歯面あるいは歯元起点を判断した。その結果を表９に示す。
次に、この傘歯車のピッチ部と歯元部における肌焼深さ及び内部硬さを測定し、更に肌焼深さ内のトルースタイト組織の観察及びオーステナイト結晶粒の混粒調査を行った。それらの結果も表９に示した。肌焼深さはマイクロビッカース硬度計（型式；ＭＶＫ−Ｅ、メーカー名；ＡＫＡＳＨＩ社製）を用いて荷重３００ｇで測定した。また、内部硬さは、ビッカース硬度計（型式；ＡＶＫ−Ｃ２、メーカー名；ＡＫＡＳＨＩ社製）を用いて荷重１０ｋｇで測定した。トルースタイト組織の観察は画像解析装置（型式；ＬＵＺＥＸ−ＩＩＩＵ、メーカー名；ＮＩＲＥＣＯ社製）を用いて以下の方法でトルースタイト面積率を調べた。ナイタール腐食後、トルースタイト組織の黒い腐食部の面積率を算出した。オーステナイト結晶粒は光学顕微鏡（型式；ＢＸ６０Ｍ、メーカー名；ＯＬＹＭＰＵＳ社製）を用いて観察した。

試験結果
表８において、Ｏ鋼はＣの含有量が本発明の範囲外であり、パーライト硬さがＨｖ２６７と良好であったが、素地硬さがＨｖ１７３と高く、冷間加工性における変形抵抗も１０００ＭＰａを超えて高かった。Ｐ鋼はＭｎの含有量が本発明の範囲外であり、素地硬さはＨｖ１５９と良好であったが、パーライト硬さが高く、冷間加工性においても変形抵抗が１０００ＭＰａより低かったが、限界加工率も劣っていた。Ｓ鋼はＢ，Ｎｂ及びＴｉを含有せず、素地硬さ及びパーライト硬さに劣り、冷間加工性もＯ鋼、Ｐ鋼より限界加工率が更に劣っていた。Ｖ鋼はＣの含有量が本発明の範囲外であり、パーライト硬さがＨｖ１７１と良好であったが、素地硬さがＨｖ１７１と高く、冷間加工性における変形抵抗も１０００ＭＰａを超えて高かった。Ｗ鋼はＭｎの含有量が本発明の範囲外であり、素地硬さがＨｖ１５８と良好であったが、パーライト硬さがＨｖ３０５と高く、また、限界加工率に劣っていた。
一方、本発明鋼のＡ〜Ｈ鋼は、素地硬さ及びパーライト硬さに優れ、冷間加工性においても変形抵抗が１０００ＭＰａ以下であり、限界加工率も７０％を超えていた。特にＣ鋼は変形抵抗が小さく、Ｄ鋼は限界加工率が８０％を超え、優れた冷間加工性を示した。本発明鋼のＡ〜Ｈ鋼は、表４におけるＨｄ_１−Ｈｄ_２値及び表６における有効Ａｌ重量％が本発明の範囲内であった。
表９において、Ｉ鋼では、Ｃの含有量が本発明の範囲外であり、ピッチ部の内部硬さが低く、歯面強度の低下により低サイクル疲労強度が低くなった。Ｊ鋼では、Ｍｎの含有量が、Ｋ鋼ではＣｒの含有量が、それぞれ本発明の範囲外であり、浸炭層にトルースタイトが５％以上析出して歯面強度が低下し、低サイクル疲労強度が低くなった。Ｌ鋼では浸炭後のオーステナイト結晶粒に混粒が発生していた。これにより、歯車ひずみや低サイクル疲労強度ばらつきが生じる。Ｍ鋼では、Ｃｒの含有量が本発明の範囲を超えて高いため、歯元曲げ強度が低下し、低サイクル疲労強度が低くなった。Ｎ鋼では、Ｎの含有量が本発明の範囲を下回ったため、浸炭後のオーステナイト結晶粒に混粒が発生した。Ｏ鋼では、低サイクル疲労強度が劣っていた。Ｑ１及びＲ１鋼では、トルースタイトが大量に析出して歯面強度が低下し、低サイクル疲労強度が低くなった。Ｓ鋼ではピッチ部及び歯元部の内部硬さに優れ、トルースタイト面積率が１％以下であったが、混粒が発生していた。Ｔ及びＵ鋼ではトルースタイトが大量に析出して歯面強度が低下した。Ｖ鋼では、Ｃの含有量が本発明の範囲を超えて高く、また有効Ａｌ重量％が低いため、混粒が発生し、靱性が低下したため、歯元強度が低下し、低サイクル疲労強度が低くなった。Ｗ鋼では、Ｃｒの含有量が本発明の範囲より低く、Ｈｄ_１−Ｈｄ_２値が本発明の範囲外であるため、トルースタイトが大量に析出して歯面強度が低下し、低サイクル疲労強度が低くなった。
Ｉ鋼はＣの含有量が本発明の範囲外である以外は、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲内であった。Ｊ鋼はＭｎの含有量及びＨｄ_１−Ｈｄ_２値が本発明の範囲外である以外は、有効Ａｌ重量％が本発明の範囲内であった。Ｋ鋼はＭｏの含有量が本発明の範囲外である以外は、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲内であった。Ｌ鋼はＣｒ，Ａｌの含有量、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲外であった。Ｍ鋼はＣｒの含有量が本発明の範囲外である以外は、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲内であった。Ｎ鋼はＡｌ及びＮの含有量が本発明の範囲外である以外は、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲内であった。Ｑ１鋼はＣｒ，Ｎｂ及びＴｉの含有量並びにＨｄ_１−Ｈｄ_２値が本発明の範囲外である以外は、有効Ａｌ重量％が本発明の範囲内であった。Ｒ１鋼はＣｒ，Ａｌの含有量、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲外であった。Ｓ鋼はＣｒ，Ａｌ，Ｎ，Ｂ，Ｎｂ及びＴｉの含有量、Ｈｄ_１−Ｈｄ_２値及び有効Ａｌ重量％が本発明の範囲外であった。Ｔ鋼は有効Ａｌ重量％が、Ｕ鋼はＨｄ_１−Ｈｄ_２値が本発明の範囲外であった。Ｖ鋼はＣ，Ａｌの含有量及び有効Ａｌ重量％が本発明の範囲外であった。Ｗ鋼はＣｒ，Ｍｎの含有量及びＨｄ_１−Ｈｄ_２値が本発明の範囲外であった。
一方、本発明鋼であるＡ〜Ｈ鋼ではトルースタイト面積率が小さく、またオーステナイト結晶粒も整粒であった。３００回強度も８０ＫＮを超え、優れた低サイクル疲労強度を示した。
尚、Ｑ１〜３３鋼及びＲ１〜２０鋼は、それぞれ特開平１１−７１６５４号公報及び特開平１０−１５２７４６号公報に記載されているものであるが、Ｑ２〜５，７，９〜１１，１３〜１５，２０，２１，２３，３２鋼及びＲ１〜１１，１３，１４，１６〜１９鋼は表４及び５におけるＨｄ_１−Ｈｄ_２値及び表６及び７における有効Ａｌ重量％が本発明の範囲外である。Ｑ６，８，１２，１６〜１９，２５，２８鋼は有効Ａｌ重量％が０．１を超えるが、Ｈｄ_１−Ｈｄ_２値が負の値をとる。また、Ｑ２２，２４，２６，２７鋼及びＲ１２，１５，２０鋼はＨｄ_１−Ｈｄ_２値が正の値をとるが、有効Ａｌ重量％が０．１未満である。
尚、本発明においては、上記実施例に限定されるものではなく、目的、用途に応じて本発明の範囲内で種々の実施例とすることができる。
【図面の簡単な説明】
図１は、歯車の低サイクル疲労破壊の形態を示す説明図である。
図２は、傘歯車を浸炭焼入焼戻するヒートパターンを示す説明図である。
図３は、傘歯車の低サイクル疲労試験を示す説明概略図である。Technical field
The present invention relates to a carburizing steel and a carburizing gear using the same, and more particularly, to a carburizing steel excellent in cold forging formability and a carburizing gear excellent in low cycle fatigue strength.
Background art
2. Description of the Related Art Conventionally, gears used for automobiles, industrial machines, and the like are made of alloy steels for machine structure such as JIS SMnC, SCr, SCM, SNCM, etc., which are subjected to hot forging, warm forging, or cold forging. After being processed and molded, it is subjected to a surface hardening treatment (such as carburizing quenching, induction quenching, nitrocarburizing treatment, etc.) in order to improve wear resistance and fatigue strength.
Japanese Patent Publication No. 7-100840 discloses a gear using steel to which a large amount of alloying element such as Mo is added, and thereby has a low cycle fatigue strength (tooth surface fatigue strength such as spalling fracture and impact tooth root). It is known that the fatigue strength is improved. However, in addition to the addition of expensive alloying elements, there is a drawback that the cost for cold forging and the machinability are deteriorated, so that the cost is greatly increased. Therefore, as shown in JP-A-10-152746 and JP-A-11-71654, a small amount of B is added in order to reduce the alloy component added to the material and to compensate for the decrease in hardenability due to this. Carburizing steels that have both workability and strength have been proposed.
However, the inventions described in these two publications do not sufficiently study the prevention of the formation of troostite in the gear, particularly in the carburized layer, and the stability of austenite crystal grains during carburization, and fail to obtain excellent low cycle fatigue strength. .
Here, the troostite means an incompletely quenched layer observed on the gear surface when gas carburizing or the like is normally performed (oxygen in a carburizing gas atmosphere diffuses from the gear surface, and Si, Mn, Unlike oxides formed with alloying elements such as Cr, the solid solution alloying elements such as Si, Mn, and Cr are deficient in the surrounding area, resulting in reduced hardenability. It refers to fine pearlite which mainly precipitates along the austenite grain boundaries of the carburized layer during cooling, and when present in a large amount, is a major factor in lowering low cycle fatigue strength.
The present invention has been made in view of the above circumstances, and an object of the present invention is to provide a carburizing steel excellent in cold forging formability and a carburized gear excellent in low cycle fatigue strength.
Disclosure of the invention
The present inventors have studied carburizing steel having excellent cold forging formability and carburized gear having excellent low cycle fatigue strength, and as a result, have completed the present invention.
That is, the carburizing steel of the present invention contains C: 0.10 to 0.30% by weight, Si: 0.50% by weight or less, Mn: 0.50 to 1.50% by weight, and P: 0.030% by weight. Hereinafter, S: 0.030% by weight or less, Cr: 0.85 to 2.00% by weight, Mo: 0.35% by weight or less, B: 0.0010 to 0.0050% by weight, Al: 0.11% 0.30% by weight, N: 0.0080 to 0.0250% by weight, Nb: 0.01 to 0.10% by weight, Ti: 0.01 to 0.10% by weight, the balance being Fe and inevitable Consisting of impurities,
Hd = 83C wt% + 5.5Mn wt% + 4.0Cr wt% + 10.5Mo wt% + 12 In the relational expression, Hd ≧ 60 ° C wt% + 12.5, and effective Al wt% = Al wt% −2. (N wt% −0.30 Ti wt%) ≧ 0.1.
C improves the case hardening depth and internal hardness after carburizing and quenching, and greatly affects the improvement of low cycle fatigue strength. However, when the content of C is less than 0.10% by weight, a long carburizing time is required to obtain the case hardening depth, so that the cost is high, and the internal strength is also reduced, and the low cycle fatigue strength of the gear is greatly reduced. I do. On the other hand, when the content exceeds 0.30% by weight, the internal hardness after carburizing and quenching increases, and the toughness of the gear teeth decreases significantly, so that the low cycle fatigue strength decreases. In addition, the cold workability and machinability, which are the manufacturing processes of the gears, are deteriorated, and the life of dies, tools, and the like is deteriorated. The lower limit of the content of C is preferably 0.13% by weight, more preferably 0.19% by weight, and the upper limit of the content of C is preferably 0.27% by weight, more preferably 0.1% by weight. 24% by weight.
Although Si is an element added as a deoxidizing agent, it is an element that forms a solid solution in ferrite and strengthens it, and particularly affects cold workability. If the content of Si exceeds 0.50% by weight, carburization may be impaired. The content of Si is preferably 0.35% by weight or less, more preferably 0.15% by weight or less. The lower limit is usually 0.03% by weight.
Although Mn is an element added as a deoxidizing agent, in order to improve the hardenability of the material, the case hardening depth and internal hardness after carburizing and quenching are improved, which greatly affects the low cycle fatigue strength. I do. However, when the Mn content is less than 0.50% by weight, the formation of troostite in the carburized diffusion layer becomes remarkable, the internal hardness is also reduced, and the low cycle fatigue strength of the gear is significantly reduced. On the other hand, if it exceeds 1.50% by weight, the transformation temperature of pearlite during soft annealing decreases, and the hardness of the pearlite portion exceeds Hv300, so that the cold workability and machinability of the gear deteriorate, and the die Since the life of tools and tools is deteriorated, the cost is significantly increased. The lower limit of the Mn content is preferably 0.8% by weight, more preferably 1.0% by weight, and the upper limit of the Mn content is preferably 1.5% by weight, more preferably 1.% by weight. 4% by weight.
P is an element that forms a solid solution in ferrite and strengthens it, and particularly an element that deteriorates cold workability. If the P content exceeds 0.5% by weight, segregation occurs at austenite grain boundaries during carburization, and the grain boundary strength of the carburized layer decreases. The content of P is preferably 0.015% by weight or less, more preferably 0.012% by weight or less. The lower limit is usually 0.002% by weight.
When S is contained in a large amount, it segregates at the austenite grain boundary during carburization, and lowers the grain boundary strength of the carburized layer. Further, it is an element that lowers the cold workability to form MnS, which is a compound with Mn, but is also an element that improves the machinability of the gear. The lower limit of the content of S is preferably 0.005% by weight, more preferably 0.008% by weight, and the upper limit of the content of S is preferably 0.020% by weight, more preferably 0.015% by weight. %.
The Cr improves the hardenability of the material, and thus improves the case hardening depth and internal hardness after carburizing and quenching, and greatly affects the improvement in low cycle fatigue strength. In particular, in steel to which B is added, troostite, which lowers low cycle fatigue strength, is likely to be generated in the carburized diffusion layer, but has a remarkable effect in preventing the formation. If the Cr content is less than 0.85% by weight, the formation of troostite in the carburized layer becomes remarkable, while if it exceeds 2.00% by weight, Cr carbide precipitates at the austenite grain boundaries during carburization, and the grain boundaries of the carburized layer are formed. The strength is reduced, which is not preferable. The lower limit of the Cr content is preferably 0.9% by weight, more preferably 1.0% by weight, and the upper limit of the Cr content is preferably 1.5% by weight, more preferably 1.3% by weight. %.
Mo improves the hardenability of the material, improves the case hardening depth and internal hardness after carburizing and quenching, and is also effective in improving low cycle fatigue strength. However, if the Mo content exceeds 0.35% by weight, the hardness increases significantly during the production of the gear, the cold workability and machinability deteriorate, and the life of the molds and tools is shortened, The cost is high. In addition, in a steel to which B is added, when a large amount of Mo is contained, troostite is easily generated in the carburized layer. The content of Mo is preferably 0.25% by weight or less, more preferably 0.20% by weight or less. The lower limit is usually 0.005% by weight.
B is precipitated at the austenite grain boundaries of the carburized layer to improve the grain boundary strength of the carburized layer, and to improve the internal hardenability after carburizing and quenching without increasing the material hardness during cold forging. Is an important element. However, when the content exceeds 0.0050% by weight because the effect of improving the hardenability of the carburized layer is small, it becomes easy to generate troostite in the carburized layer, and further, the effect of improving the hardenability is saturated. On the other hand, hot forgeability or cold forgeability deteriorates, and if it is less than 0.0010% by weight, the internal hardenability during carburizing and quenching decreases, which is not preferable. The content of B is preferably 0.0010 to 0.0035% by weight, more preferably 0.0015 to 0.0030% by weight.
Al is an element added as a deoxidizing agent, and has an effect of reacting with N in steel to form AlN and preventing coarsening of austenite crystal grains during carburizing heating. When the effective Al weight% = Al weight% -2 (N weight% −0.30 Ti weight%) ≧ 0.1, if the Al content is less than 0.11% by weight, troostite is formed in the carburized layer. In the case of a gear having a high cold working ratio, abnormal growth of austenite crystal grains during carburization is observed. When the Al content exceeds 0.3% by weight, the surface of the gear is incomplete during carburization. The quenched layer is undesirably deep. If the effective Al weight% is less than 0.1% by weight, the hardenability decreases and the troostite of the carburized layer becomes remarkable. The content of Al is preferably 0.11 to 0.20% by weight, and more preferably 0.11 to 0.15% by weight.
Further, the content of AlN formed is preferably 200 to 600 ppm, more preferably 250 to 550 ppm, and still more preferably 300 to 500 ppm.
The N reacts with Al to precipitate AlN, but if the content of N is less than 0.0080% by weight, in the cold-worked gear, the amount of precipitated AlN is reduced and austenite crystal grains during carburization are reduced. Coarse. On the other hand, if it exceeds 0.0250% by weight, the amount of solid solution B decreases, and the hardenability decreases. The lower limit of the N content is preferably 0.008% by weight, more preferably 0.012% by weight, and the upper limit of the N content is preferably 0.025% by weight, more preferably 0.020% by weight. is there.
The Nb reacts with N or C to generate Nb (C, N) and has an effect of preventing austenite crystal grains from being coarsened during carburizing heating. If the Nb content is less than 0.01% by weight, the above effect is not exhibited, and if it exceeds 0.10% by weight, the effect of preventing austenite crystal grains from being coarsened during carburizing heating is saturated, resulting in high cost. Is not preferred. The content of Nb is preferably 0.01 to 0.07% by weight, more preferably 0.01 to 0.05% by weight.
The Ti reacts with N to generate TiN, which has an effect of preventing austenite crystal grains from being coarsened during carburizing heating. If the content of Ti is less than 0.01% by weight, the above effect is not exhibited. If the content exceeds 0.10% by weight, the effect of preventing austenite crystal grains from being coarsened during carburizing heating is saturated, and the cost is increased. Is not preferred. The content of Ti is preferably 0.01 to 0.07% by weight, and more preferably 0.01 to 0.055% by weight.
The above formula Hd = 83C wt% + 5.5Mn wt% + 4.0Cr wt% + 10.5Mo wt% + 12 is a formula for obtaining a parameter representing the magnitude of the internal hardness after carburizing and quenching of the carburizing steel having the above composition. The larger the value, the higher the internal hardness after carburizing and quenching.
When the Hd thus obtained satisfies the above formula, Hd ≧ 60 ° C. wt% + 12.5, the martensite ratio can be increased to 90% or more after carburizing, quenching and tempering.
Further, in the state of the material which has been appropriately soft-annealed when the above-described carburizing steel is processed into a predetermined shape before carburizing and quenching, the hardness is Hv 170 or less (preferably Hv 165 or less, more preferably Hv 160 or less. ) And a hardness of Hv300 or less (preferably Hv295 or less, more preferably Hv290 or less) at a micro Vickers hardness of 10 g under a load. This improves the cold forgeability of the carburizing steel.
The carburized gear of the present invention is manufactured using the carburizing steel described above, has a case hardening depth Hv513 of 0.5 mm or more from the surface, and has a troostite area ratio within the case hardening depth. It is not more than 5%.
The method of processing the carburized gear is not particularly limited, but usually, any one of hot forging, warm forging, and cold forging is performed, followed by machining and carburizing and quenching. In this case, the hardness of the surface layer is usually represented by Vickers hardness, and the case-hardening depth Hv513 is preferably 0.6 mm or more from the surface, more preferably 0.65 mm or more, and still more preferably 0 mm or more. 0.7 mm or more. However, the upper limit is usually 1.5 mm, more preferably 1.2 mm. If the above case hardening depth is less than 0.5 mm, it will be broken starting from the inside of the gear, which is not preferable. On the other hand, if the case hardening depth is 0.5 mm or more, the fatigue strength is remarkably excellent. The method for measuring the hardness of the gear will be described in Examples.
The fracture due to low cycle fatigue of the above gears is caused by tooth surface fatigue peeling caused by plastic deformation inside the tooth surface due to high contact stress on the tooth surface, and the inside of the tooth root due to high bending stress on the tooth root. There is a root impact fatigue fracture caused by plasticity and an increase in root surface stress (see FIG. 1). Generally, a carburized and quenched gear has a gentle C concentration distribution from the surface C concentration of 0.8% toward the inside of the tooth (uncarburized portion). Since the strength and toughness of the (uncarburized part) are important, in addition to the suppression of the formation of troostite in the carburized layer and the stability of the austenitic crystal grains caused by the toughness, the quenched structure inside the tooth (martenite) caused by the internal strength Site rate) is important.
The “area fraction of troostite” refers to an average area fraction in a region from the surface after carburizing and quenching to a depth of 0.5 mm, preferably 4.7% or less, more preferably 4. 5% or less. If it exceeds 5%, the fatigue strength is reduced, and there is a fear that peeling may occur on the tooth surface of the gear (the portion where the teeth contact each other when the gear meshes).
The internal hardness of the pitch portion of the carburized gear may be Hv350 or more, preferably Hv370 or more, more preferably Hv390 or more. If it is less than Hv350, the tooth surface fatigue strength is undesirably reduced.
Further, the internal hardness of the root portion of the carburized gear can be Hv300 or more, preferably Hv330 or more, more preferably Hv360 or more. If it is less than 300, the tooth root impact fatigue strength is undesirably reduced. In addition, the above-mentioned "root part" means an R part which goes from the pitch circle part of the tooth toward the tooth bottom.
As a preferable combination of the internal hardness of the pitch portion and the root portion, the pitch portion is Hv350 or more, and the root portion is Hv300 or more, more preferably, the pitch portion is Hv370 or more, and the root portion is Hv330 or more. Preferably, the pitch portion is Hv390 or more and the root portion is Hv360 or more.
Further, as a preferable combination of the internal hardness of the pitch portion, the internal hardness of the root portion, and the area ratio of the troostite, (1) the hardness of the pitch portion is Hv350 or more, and the hardness of the root portion is Hv300 or more. And the troostite area ratio is 5% or less, (2) more preferably, the hardness of the pitch portion is Hv370 or more, and the hardness of the tooth root portion is Hv330 or more, and the troostite area ratio is 4.8% or less, (3) More preferably, the hardness of the pitch portion is Hv390 or more, the hardness of the root portion is Hv360 or more, and the troostite area ratio is 4.7% or less. (4) Particularly preferably, the hardness of the pitch portion is Hv400 or more, the root hardness is Hv390 or more, and the troostite area ratio is 4.5% or less.
In the case of manufacturing a carburized gear using the above-mentioned carburizing steel, a steel having the above structure and properties can be obtained by performing normal carburizing, quenching and tempering. In particular, the structure is affected by the mass of the gear and the carburizing and quenching medium, and the mass of the gear is equivalent to a round bar, and the outer diameter is preferably 40 mm or less, more preferably 30 mm or less, and still more preferably 25 mm or less. The quenching medium is preferably oil quenching at 200 ° C. or lower. However, the carburizing atmosphere may be performed in any atmosphere such as gas carburizing, carburizing and nitriding, and vacuum carburizing.
According to the carburizing steel of the present invention, it has a specific elemental composition, and in the relational expression of Hd = 83C weight% + 5.5Mn weight% + 4.0Cr weight% + 10.5Mo weight% + 12, Hd ≧ 60 ° C. weight. By setting it to% + 12.5, a carburized and quenched structure having a martensite ratio of 90% or more can be obtained, and abnormal growth of austenite crystal grains can be prevented. Further, when processing into a predetermined shape before carburizing and quenching, by performing appropriate softening annealing, cold forging formability can be improved and processing can be easily performed. Therefore, gears and shafts used for automobiles, industrial machines, and the like can be easily manufactured using the carburizing steel. Further, according to the carburized gear of the present invention, since the stress concentration at the root portion, which is the effect of the cold forged gear, and the forging line along the tooth profile can be obtained, the strength is higher and the torque is higher. It is excellent in low cycle fatigue strength (tooth surface fatigue strength and tooth root fatigue strength) which is the resistance of the region.
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be specifically described with reference to examples.
1. Production and evaluation of carburizing steel
(1) Hardness and effective Al weight% obtained from the composition of carburizing steel
Steels A to S having the chemical compositions shown in Tables 1 to 3 (A to H were steels of the present invention, and I to W were comparative steels. Among the comparative steels, Q1 to 33 steel and R1 to 20 steel were No. 152746 and Japanese Unexamined Patent Application Publication No. 11-71654. In addition, S steel is equivalent to JIS standard SNCM420 widely used for gears requiring low cycle fatigue strength.) 30 kg was melted in a vacuum melting furnace.
Tables 4 and 5 show the Hd calculated from each component amount. ₁ = 83C wt% + 5.5Mn wt% + 4.0Cr wt% + 10.5Mo wt% + 12 and Hd ₂ = 60 ° C wt% + 12.5 and their difference Hd ₁ -Hd ₂ showed that. Tables 6 and 7 show the calculation results of the effective Al weight%. In Tables 4 to 7, those having a component element content outside the range of the present invention are indicated by an asterisk beside the numerical value.

(2) Production and evaluation of carburizing steel
The ingots A to P, Q1, R1, and S were heated at 1200 ° C. or higher for 0.5 hour, and then hot forged at a temperature of 1000 to 1200 ° C. to produce a round bar having a diameter of 30 mm. This was annealed from 900 ° C. to 600 ° C. at a cooling rate of 75 ° C./min to obtain a carburizing steel.
Among the carburizing steels obtained above, the carburizing round bars of A to H, O, P and S have hardness (hardness of the base and pearlite) and cold workability (70% upsetting). (Deformation resistance and critical working ratio) were measured by the following methods. Table 8 shows the results.
The substrate hardness was measured at a load of 10 kg using a Vickers hardness tester (model: AVK-C2, manufacturer: manufactured by AKASHI). The hardness of the pearlite portion was measured by measuring the micro Vickers hardness under a load of 10 g.
The deformation resistance was determined by applying a test piece (without cutout) having a diameter of 10 mm and a height of 15 mm using a 100 t universal testing device (model: RH-100, manufacturer name; manufactured by Shimadzu Corporation) at a load cell moving speed of 1 mm / min. Table 8 shows the deformation resistance obtained by measuring the compression load at the time of% upsetting and using the deformation resistance measurement method by the end face constrained compression proposed by the Japan Society for Plastic Working. 1980) pp. 529-532).
The limit processing rate is determined by performing an end face restraint test on the test piece using the above-described apparatus, applying a load cell moving speed of 1 mm / min, and setting the upsetting rate when a crack occurs at the end (circumferential portion). (See Plasticity and Processing No. 241, (1981) pp. 139 to 144, published by the Japan Society for Technology of Plasticity).

2. Manufacture and evaluation of carburized gears
Of the carburizing steels obtained above, for A to N, Q1, R1, and SW, the carburizing round bar was cold forged with a module of 4.8, a number of teeth of 10, and a pitch circle diameter of 48. A 0.8 mm bevel gear was manufactured by carburizing, quenching and tempering with the heat pattern shown in FIG. This bevel gear was set in a hydraulic fatigue test apparatus as shown in FIG. 3, and a bevel gear fatigue test was performed in the following manner. A jig having a curvature equivalent to that of an actual mating gear was prepared for the test gear, and a complete one-sided fatigue test was performed. Using the acoustic emission, the crack generation time was regarded as the life, and the tooth surface or the starting point of the tooth root was determined by magnetic particle inspection. Table 9 shows the results.
Next, the case hardening depth and the internal hardness at the pitch part and the tooth root part of this bevel gear were measured, and further, the troostite structure within the case hardening depth was observed and the austenite crystal grains were mixed. The results are also shown in Table 9. The case hardening depth was measured with a load of 300 g using a micro Vickers hardness tester (model: MVK-E, manufacturer name: manufactured by AKASHI). The internal hardness was measured at a load of 10 kg using a Vickers hardness tester (model: AVK-C2, manufacturer: manufactured by AKASHI). For the observation of the troostite structure, the area ratio of troostite was determined by the following method using an image analyzer (model: LUZEX-IIIU, manufacturer name: NIRECO). After the nital corrosion, the area ratio of the black corroded portion of the troostite structure was calculated. The austenite crystal grains were observed using an optical microscope (model: BX60M, manufacturer name: manufactured by OLYMPUS).

Test results
In Table 8, the O steel had a C content outside the range of the present invention and had a good pearlite hardness of Hv267, but had a high base hardness of Hv173 and a deformation resistance of 1000 MPa in cold workability. It was higher than it was. The P steel had a Mn content outside the range of the present invention and had a good base hardness of Hv159, but had a high pearlite hardness and a low deformation resistance in cold workability of less than 1000 MPa. The processing rate was also poor. The S steel did not contain B, Nb and Ti, was inferior in the base hardness and the pearlite hardness, and was further inferior in the cold workability to the critical work ratio than the O steel and the P steel. The V steel had a C content outside the range of the present invention and had a good pearlite hardness of Hv171, but had a high base hardness of Hv171 and a high deformation resistance in cold workability exceeding 1000 MPa. . The W steel had a Mn content outside the range of the present invention and had a good base hardness of Hv158, but had a high pearlite hardness of Hv305 and was inferior to the critical working ratio.
On the other hand, the A to H steels of the present invention steels were excellent in the base hardness and the pearlite hardness, the deformation resistance was 1000 MPa or less even in the cold workability, and the critical working ratio exceeded 70%. In particular, the C steel has a low deformation resistance, and the D steel has a critical working ratio exceeding 80%, and has excellent cold workability. The steels A to H of the present invention steels are listed as Hd ₁ -Hd ₂ The values and effective Al wt% in Table 6 were within the scope of the present invention.
In Table 9, in Steel I, the content of C was out of the range of the present invention, the internal hardness of the pitch portion was low, and the low cycle fatigue strength was low due to the decrease in tooth surface strength. The content of Mn in the J steel and the content of Cr in the K steel are out of the range of the present invention, respectively, and 5% or more of troostite precipitates in the carburized layer to lower the tooth surface strength, and low cycle fatigue. Strength decreased. In the L steel, mixed grains were generated in the austenite crystal grains after carburization. This causes gear strain and low cycle fatigue strength variation. In the M steel, since the Cr content was higher than the range of the present invention, the root flexural strength decreased and the low cycle fatigue strength decreased. In N steel, since the N content was less than the range of the present invention, austenite crystal grains after carburization were mixed. In O steel, the low cycle fatigue strength was inferior. In the Q1 and R1 steels, a large amount of troostite was precipitated, the tooth surface strength was reduced, and the low cycle fatigue strength was low. In the case of S steel, the internal hardness of the pitch portion and the root portion was excellent, and the area ratio of troostite was 1% or less, but mixed grains occurred. In T and U steels, a large amount of troostite was precipitated, and the tooth surface strength was reduced. In steel V, since the content of C is higher than the range of the present invention and the effective Al weight% is low, mixed grains are generated and the toughness is reduced, so that the root strength is reduced and the low cycle fatigue strength is reduced. Got lower. In W steel, the Cr content is lower than the range of the present invention, and Hd ₁ -Hd ₂ Since the value was outside the range of the present invention, a large amount of troostite was precipitated, the tooth surface strength was reduced, and the low cycle fatigue strength was low.
Steel I has Hd except that the content of C is outside the scope of the present invention. ₁ -Hd ₂ The values and effective Al wt% were within the scope of the present invention. J steel has Mn content and Hd ₁ -Hd ₂ Except for values outside the scope of the invention, the effective Al wt% was within the scope of the invention. K steel has the same structure as that of Hd except that the content of Mo is out of the range of the present invention. ₁ -Hd ₂ The values and effective Al wt% were within the scope of the present invention. L steel has Cr and Al contents, Hd ₁ -Hd ₂ The values and effective Al wt% were outside the scope of the present invention. M steel is Hd except that the Cr content is outside the scope of the present invention. ₁ -Hd ₂ The values and effective Al wt% were within the scope of the present invention. N steel is Hd except that the contents of Al and N are outside the scope of the present invention. ₁ -Hd ₂ The values and effective Al wt% were within the scope of the present invention. Q1 steel has Cr, Nb and Ti contents and Hd ₁ -Hd ₂ Except for values outside the scope of the invention, the effective Al wt% was within the scope of the invention. R1 steel has Cr and Al contents, Hd ₁ -Hd ₂ The values and effective Al wt% were outside the scope of the present invention. S steel has Cr, Al, N, B, Nb and Ti content, Hd ₁ -Hd ₂ The values and effective Al wt% were outside the scope of the present invention. Effective Al weight% for T steel, Hd for U steel ₁ -Hd ₂ The value was outside the range of the present invention. The steel V was out of the scope of the present invention in the content of C and Al and the effective Al weight%. W steel has Cr and Mn contents and Hd ₁ -Hd ₂ The value was outside the range of the present invention.
On the other hand, in the A to H steels of the present invention, the troostite area ratio was small, and the austenite crystal grains were also sized. The 300-time strength also exceeded 80 KN, showing excellent low cycle fatigue strength.
The Q1 to 33 steels and the R1 to 20 steels are described in JP-A-11-71654 and JP-A-10-152746, respectively, but Q2 to 5, 7, 9 to 11, 13 １５15,20,21,23,32 steel and R1-11,13,14,16-19 steel are Hd in Tables 4 and 5. ₁ -Hd ₂ The values and the effective Al wt% in Tables 6 and 7 are outside the scope of the present invention. Q6,8,12,16-19,25,28 steel has an effective Al weight% exceeding 0.1, ₁ -Hd ₂ The value is negative. In addition, Q22, 24, 26, 27 steel and R12, 15, 20 steel are Hd. ₁ -Hd ₂ The value takes a positive value, but the effective Al weight% is less than 0.1.
It should be noted that the present invention is not limited to the above embodiments, and various embodiments can be made within the scope of the present invention depending on the purpose and application.
[Brief description of the drawings]
FIG. 1 is an explanatory diagram showing a form of low cycle fatigue fracture of a gear.
FIG. 2 is an explanatory diagram showing a heat pattern for carburizing, quenching and tempering a bevel gear.
FIG. 3 is an explanatory schematic diagram showing a low cycle fatigue test of a bevel gear.

Claims

C: 0.10 to 0.30% by weight, Si: 0.50% by weight or less, Mn: 0.50 to 1.50% by weight, P: 0.030% by weight or less, S: 0.030% by weight or less , Cr: 0.85 to 2.00% by weight, Mo: 0.35% by weight or less, B: 0.0010 to 0.0050% by weight, Al: 0.11 to 0.30% by weight, N: 0. 0080-0.0250% by weight, Nb: 0.01-0.10% by weight, Ti: 0.01-0.10% by weight, the balance being Fe and unavoidable impurities,
Hd (hardness) = 83C weight% + 5.5 Mn weight% + 4.0Cr weight% + 10.5Mo weight% + 12, Hd ≧ 60 ° C weight% + 12.5, and effective Al weight% = Al A carburizing steel, wherein wt%-2 (N wt%-0.30 Ti wt%) ≥ 0.1.

A case hardening depth of 0.5 mm or more from the surface produced by using the carburizing steel according to claim 1 and having a hardness of Hv 513 by carburizing and quenching and tempering, and a troostite within the case hardening depth. A carburized gear having an area ratio of 5% or less.

The carburized gear according to claim 2, wherein the internal hardness of the pitch portion of the carburized gear is Hv350 or more, and the internal hardness of the root portion is Hv300 or more.