JP3580441B2

JP3580441B2 - Ni-base super heat-resistant alloy

Info

Publication number: JP3580441B2
Application number: JP11016695A
Authority: JP
Inventors: 光司佐藤; 誠小森; 庸田村; 光浩安藤; 善裕加田; 丈博大野
Original assignee: Hitachi Metals Ltd
Current assignee: Proterial Ltd
Priority date: 1994-07-19
Filing date: 1995-04-11
Publication date: 2004-10-20
Anticipated expiration: 2019-10-20
Also published as: JPH0885838A; DE19525983A1

Description

【０００１】
【産業上の利用分野】
本発明は、主に９００℃以下の温間、熱間の温度域で使用される温間、熱間鍛造金型や、ＣｕやＡｌまたはそれらの合金の熱間押出し工具や部品等に使用されるＮｉ基超耐熱合金に関するものである。
【０００２】
【従来の技術】
近年の自動車産業の部品共通化と作業工数低減の流れを受けて、自動車用のクランクシャフトやコンロッド等自動車部品を成形する温・熱間鍛造金型の寿命向上が望まれている。しかし、従来のＳＫＤ６１に代表される熱間工具鋼では、そのニーズに応えきれていないのが現状である。また、ＣｕやＡｌまたはそれらの合金の熱間押し出し工具には、Ａ２８６のようなＦｅ基超耐熱合金が用いられたり、インコネル７１８合金が使用されたりしているが、この用途においても寿命改善による工数の低減が求められている。
これらの他に、継ぎ目無し鋼管の製造に用いる穿孔プラグやマンドレル等、１２００℃程度の高温域で使用される熱間工具を対象として特開平３−６１３４５号、また、１０００℃〜１１５０℃の高温大気雰囲気中で使用される恒温鍛造金型を対象として特開平４−４１６４１号等のＮｉ基超耐熱合金が開示されている。
さらに、本発明と目的は異なるが、Ｎｉ−Ｃｒ−Ｗを基本とする超耐熱合金としては、例えば特公昭５４−３３２１２号、特公昭６０−５８７７３号、特公昭５８−５０２号、特公昭５６２−１０６１号に記載の合金等が知られている。
【０００３】
【発明が解決しようとする課題】
熱間や温間で用いる鍛造金型、あるいはＣｕやＡｌまたはそれらの合金の押し出し工具には、熱間加工が可能で、被削性にも優れ、さらに常温での高い耐力と靭性、および７００℃〜９００℃程度の高温域での優れた熱間摺動性と熱間での耐摩耗性および耐ヒートクラック性とが要求される。
ここで、熱間摺動性および熱間での耐摩耗性の改善には、酸化皮膜の密着性と安定性および高温耐力を高めることが要求される。しかし、従来のＳＫＤ６１クラスの熱間工具鋼では、７００℃を越えると焼き戻しマルテンサイトの軟化が生じて高温耐力が低下し、同時に鋼中のＣｒ量が低いために密着性および安定性に優れた酸化皮膜が形成されず、被加工材と工具が直接接触するため早期に摩耗を生じ、最後には焼付きを生じてしまう。
【０００４】
一方、耐ヒートクラック性の改善には、熱膨張係数の低減と高い高温耐力が要求される。しかし、Ａ２８６のようなＦｅ基超耐熱合金は、熱膨張係数が高く、また析出強化相であるガンマプライム相は、Ｎｉ_３Ｔｉを主とする組成のために７００℃を越える温度域で安定相であるイータ相に変態してしまい、その結果、高温強度が低下してしまう問題がある。
また、熱間工具は被削性が優れることが要求されるが、Ａ２８６のような超耐熱合金は、一般に被削性がマルテンサイト系の材料に比べて大幅に悪い欠点がある。
ところで、Ｎｉをベースとする熱間工具用の超耐熱合金もいくつか提案されているが、例えば、特開平３−６１３４５号に開示される合金は、耐酸化性が十分でなく、熱間での耐摩耗性が不十分である。また、特開平４−４１６４１号に開示される合金は、恒温鍛造型を意図しているために、高い摺動速度下での熱間摺動性が不十分な問題があった。
【０００５】
また、前述のＮｉ−Ｃｒ−Ｗ系統のＮｉ基超耐熱合金は、いずれも１０００℃付近の非常に高い温度域で、耐熱性と高温での大きいクリープ強度を第一目的として使用することを前提としている。したがって、従来のＷを多量に添加するＮｉ基超耐熱合金は、いずれもＷの固溶強化により高温クリープ強度を増加させたものであり、本発明の目的とする７００〜９００℃における高強度、熱間摺動性の改善に対しては十分ではなく、本発明が意図する利用分野では特性上使用することは困難であった。
【０００６】
【課題を解決するための手段】
本発明者らはこれら熱間工具鋼とＦｅ基およびＮｉ基超耐熱合金の問題点を鑑み、熱間工具鋼の代替を主目的の一つとして熱間加工が可能で、かつ耐酸化性、高温強度、熱膨張係数、および被削性の改善が可能な材料の発明に鋭意取り組んだ。その結果、Ｗを固溶限まで含むＮｉを主とするオーステナイト相を主相とし、時効処理によって析出させるＮｉ_３（Ａｌ，Ｔｉ）を主とするガンマプライム相、Ｗの固溶体（α−Ｗ）およびＭＣ型１次炭化物の最適な含有量を見出すことで、特に熱間工具に適した熱間摺動性、熱間での耐摩耗性と耐ヒートクラック性および耐酸化性に優れたＮｉ基超耐熱合金を発明するに至った。さらに、この組織を有する合金で、高温大気中で生成する表面スケールの構造が、表面側から、Ｃｒを主とする外部酸化層、粒状α−Ｗ層、およびＡｌ，Ｔｉを主とする内部酸化層の少なくとも３層以上の構造からなるものは、酸化皮膜の密着性を改善し、高温耐酸化性と熱間摺動性の改善に役立つことを新規に見出した。
【０００７】
本発明に係るＮｉ基超耐熱合金の特徴は、下記に示される第１発明や第２発明の構成要件から理解できるように、合金組成のみに存在するのではない。すなわち、本発明のＮｉ基超耐熱合金と、従来のＷ含有のＮｉ基超耐熱合金の決定的な相違点は、本発明ではガンマプライム相に特徴づけられる時効処理された組織を有することである。時効処理は６００〜８５０℃で行なうので、従来のＮｉ基超耐熱合金のように１０００℃付近以上で使用する用途では、せっかく析出した時効硬化相が固溶するから事前の時効処理自体が無意味であり、認識されていなかったのである。
本発明では、従来の類似の組成のＮｉ基超耐熱合金では、採用できなかった時効析出相を積極的に利用する。したがって、本発明のＮｉ基超耐熱合金の組織は、１０〜３０％の多量のガンマプライム相が存在することが特徴となっている。
【０００８】
即ち、本発明は、時効処理された組織を有し、該組織を構成する相が原子％から式（１）により計算されるＭＣ型炭化物量で０．０２〜１．５％、原子％から式（２）により計算されるガンマプライム量で１０〜３０％、および体積％で表されるα−Ｗ量で０．５〜３０％と、残部がＮｉを主体とするオーステナイト相からなることを特徴とするＮｉ基超耐熱合金である。
（１）２[Ｃ]
（２）４（０．０２６〔Ｃｒ〕＋０．１３〔Ｍｏ〕＋０．１３〔Ｗ〕＋０．６１〔Ａｌ〕＋０．６８〔Ｔｉ〕＋０．５〔Ｎｂ〕＋０．５〔Ｔａ〕−〔Ｃ〕）
（但し、上記式のうち、無添加の元素は０として計算する）
【０００９】
上記本発明を具現化するためのＮｉ基超耐熱合金の化学組成は、重量％で、Ｃ０．００２〜０．１５％、Ｓｉ２％以下、Ｍｎ３％以下、Ｃｒ１０％を越え２５％以下、Ｗ１０〜３０％、Ｆｅ１５％以下、Ａｌ０．４〜２．５％、Ｔｉ０．４〜３．５％を含み、残部が不可避の不純物とＮｉからなる。
【００１０】
上記の本発明の望ましい合金の組成は、さらにＮｉの一部を置換して、重量％で、３．０％以下のＮｂと３．０％以下のＴａの１種または２種を添加できる。またさらに上記の組成に加えてＮｉの一部を置換して、重量％で、１０％以下のＭｏを、Ｗ＋２Ｍｏ≦３０の範囲で添加すること、さらに重量％で、０．２％以下のＺｒと０．０２％以下のＢの１種または２種を添加すること、重量％で、０．０２％以下のＭｇと０．０２％以下のＣａの１種または２種を添加することもできる。以上の合金元素のうち、Ｎｂ以下に記述されるものは、いずれも可能なすべての組合せで選択的に添加できるものである。また、選択的に添加された元素は、前述のガンマプライム量の推定式を満足する必要がある。
【００１１】
【作用】
本発明のＮｉ基超耐熱合金で規定する組織の作用について述べる。
ＭＣ型の一次炭化物は、本発明のＮｉ基超耐熱合金中において最も硬質の相であり、本合金製造の熱間加工中にオーステナイト結晶粒の粗大化を抑制する作用と、使用中の熱間摺動性を高める作用を持つ。ＭＣ炭化物量は、添加したＣの全量がＴｉ，ＮｂおよびＴａと結びついて存在しているものと考えると、原子％で表されるＣ量の２倍がＭＣ型炭化物の含有量とみなすことができる。したがって、前述のＭＣ型炭化物量は（１）式の２［Ｃ］で計算される値とする。このＭＣ型炭化物の効果を発揮するために、ＭＣ量は原子％で最低０．０２％を越える添加が必要であるが、１．５％を越える過度の添加は、ＭＣ型炭化物の連鎖状組織を形成し、その部分がヒートクラック発生の起点となり、工具寿命を低下させる。よって原子％から計算されるＭＣ型炭化物量は、０．０２〜１．５％の添加とする。好適なＭＣ炭化物の量は０．１〜１．０％である。
【００１２】
常温や高温強度の向上には、時効処理によるガンマプライム（γ’）相の析出が最も効果を及ぼす。ガンマプライム相の基本組成は、Ｎｉ_３Ａｌで表わされるが、多元系の合金においては、実際には、種々の添加元素がガンマプライム相のＡｌサイトに固溶する。このガンマプライム相の析出量を実測するのは、大変困難であるため、本発明においては、文献値（原田広史、山崎道夫：「Ｔｉ，Ｔａ，Ｗを含むγ’析出強化型Ｎｉ基耐熱鋳造合金の合金設計」，鉄と鋼，ｖｏｌ．６５，（１９７９），ｐ１０５９）を用いて、後述の実施例１の本発明合金Ｎｏ．５のγ相とγ’相の組成および両相の割合を求め、原子％から各元素がγ’相のＡｌサイトに入る係数を求めた。ここで、本発明合金の成分範囲から、Ｎｉ_３ＡｌのＡｌサイトに固溶する元素として、Ｃｒ，Ｍｏ，Ｗ，Ｔｉ，Ｎｂ，Ｔａを考えた。ただし、実施例１のＮｏ．５は、ＭｏとＴａを含まない成分系であるため、ＭｏとＴａの係数は計算できない。そこでＭｏとＴａのγ’相のＡｌサイトを占める係数は、それぞれ同族のＷとＮｂの係数を用いることにした。
【００１３】
また、その計算に際して、Ｔｉ，Ｎｂ，Ｔａは、添加したＣの全量とＭＣ型炭化物を形成した残部がγ−γ’相の平衡に関与するものとした。この計算結果から、添加元素のうち、γ’相のＡｌサイトに入る総量を求め、さらにこの４倍を計算γ’量とした。即ち、本発明におけるガンマプライム量は、次式で計算されるものとし、前述の（２）式に相当する。
計算γ’＝４（０．０２６［Ｃｒ］＋０．１３［Ｍｏ］＋０．１３［Ｗ］＋０．６１［Ａｌ］＋０．６８［Ｔｉ］＋０．５［Ｎｂ］＋０．５［Ｔａ］−［Ｃ］）
（［］内の元素は原子％を表わし、無添加の元素は０として計算する）
【００１４】
次に本発明のＮｉ基超耐熱合金では、γ’相を析出させるための熱処理が施されているものである。この熱処理は、まず熱間加工後の組織を均一にし、析出物を一旦固溶させるための、固溶化処理を行なう。この固溶強化の温度は、９００〜１２００℃の範囲が望ましい。９００℃以下では析出物の固溶が不十分で、１２００℃以上では結晶粒が粗大化し、強度、耐ヒートクラック性に不利である。但し、常温の強度を重視する場合には、鍛造の歪を残す目的で固溶化処理をより低温で行なうか、あるいは省略することもできる。
【００１５】
引き続き時効処理を行なうが、この時効処理は、本発明合金のＮｉ基超耐熱合金では、γ’相を析出させ、高温強度を増加させるために不可欠である。時効処理は、６００〜８５０℃の温度が望ましい。６００℃未満では、γ’相の析出に多大な時間を要し、８５０℃を越えると析出するγ’相が粗大になるか、またはγ’相が固溶したままで析出せず、十分な高温強度が得られない。
この時効処理は、１回または２回以上の処理を行なうことができるが、処理時間は合計で少なくとも３Ｈｒ以上が必要である。
時効処理により析出した上記計算γ’量が１０％を下回ると、高温強度が低下し、それに伴い、熱間摩耗や、熱間摺動に対する抵抗が低下して工具寿命の低下を引き起こす。逆に、計算γ’量が３０％を越えると、高温変形抵抗が増加し、熱間鍛造が困難となる。よって上記式による計算γ’量は、１０〜３０％の範囲とする。好適には、１３〜２５％の範囲である。
【００１６】
なお、前述の特公昭５４−３３２１２号、特公昭６０−５８７７３号、特公昭５８−５０２号、特公昭５６−２１０６１号に記載の合金は、Ｗ添加による固溶強化を目的としており、ガンマプライム形成元素をほとんど含んでいない。あるいは少量含む場合でも時効処理を行なわないので、ガンマプライム相による強化は図られていない。なぜなら、使用温度が１０００℃付近であるのでγ’相は、使用中に固溶してしまうからであり、この場合のγ’形成元素であるＡｌ，Ｔｉ，Ｎｂ等の作用は、例えばＡｌは耐酸化性の向上、Ｔｉ，Ｎｂは炭化物形成による強度向上の作用等で、本発明のように積極的にγ’相による強化を狙ったものではなく目的が全く異なる。
さらに、上記の各合金は１０００℃付近の高温クリープ強度を重視するため、１２００℃以上の温度で固溶化処理を行なったままか、あるいは固溶化処理後、１０８０℃以上再加熱を行なった状態で使用するので、結晶粒が粗大化しており、本発明の目的とする７００〜９００℃における強度、耐ヒートクラック性を必要とする用途には不利である。
【００１７】
次に、均一に分散したα−Ｗ相（ｂｃｃ構造を有するＷの固溶体）は、本発明合金に必須の重要な合金相である。α−Ｗ相は、合金の熱膨張係数の低下に寄与するとともに、熱間摺動時の焼付きを防止する。さらに、α−Ｗは、以下でも述べるが、酸化皮膜の密着性を改善する効果をももたらす。このα−Ｗの効果は、画像解析などの実側による体積％にして、０．５％以上から発揮されるが、３０％を越えると、耐酸化性と、熱間加工性の低下を招くので、α−Ｗの含有量は体積％にして、０．５〜３０％の範囲とする。好適には、３〜２０％の範囲である。
本発明のＮｉ基超耐熱合金の特徴は、γ’相とα−Ｗ相が同時に存在することであり、これは従来の超耐熱合金には見られなかったことである。
すなわち、通常のγ’析出強化型超耐熱合金には当然γ’相が存在するが、α−Ｗ相は存在しない。
【００１８】
従来のＮｉ基超耐熱合金のうち、高Ｗである合金はα−Ｗ相が存在する可能性があるが、それはせいぜい結晶粒界や粒内に少量析出させてクリープ強度を高める効果を図ったもので、本発明における作用効果とは全く異なる。また、従来のこれらの合金には、前述のようにγ’相は存在しない点で決定的な相違がある。本発明合金においては、γ’相とα−Ｗ相が同時に存在することで、初めて７００〜９００℃における強度、熱間摺動、熱間摩耗、耐ヒートクラック、耐酸化性等の要求特性を同時に満足することが可能となったのである。
さらに、本発明のＮｉ基超耐熱合金は、例えば実施例１における図１に示すように高温大気中で生成する表面スケールの構造が、表面側から、Ｃｒを主とする外部酸化層、粒状α−Ｗ層、およびＡｌ，Ｔｉを主とする内部酸化層の少なくとも３層以上の構造をとることを特徴とする。この酸化皮膜形成のメカニズムは明らかではないが、酸化皮膜構造がこの少なくとも３層以上の構造を取らない場合には、酸化皮膜の密着性と耐酸化性の低下が生じるようになる。
【００１９】
例えば、本発明のＮｉ基超耐熱合金においても５μｍを越えるα−Ｗ粒子が表面に出ているとその部分はＷＯ_３の酸化皮膜が生成し、その部分は、十分な密着性を持たないが、全体としてα−Ｗの総量が３０％以下に抑えられているので、工具自身の寿命を大きく低下させるほどの要因とはならない。なお、スケール中の粒状α−Ｗ層は析出相のα−Ｗ相から変化させ、粒子径は１μｍ以下のものが多くするのが望ましい。よって本発明合金を大気中で加熱した場合の表面スケールの構造は、表面側から、Ｃｒを主とする外部酸化層、粒状α−Ｗ層、およびＡｌ，Ｔｉを主とする内部酸化層の少なくとも３層以上の構造をとることが望ましい。
【００２０】
さらに、上記の量のＭＣ型炭化物と、γ’相およびα−Ｗ相と不可避の不純物を除き、残部は、Ｎｉを主体とするオーステナイト相である。この母相は、Ｗの固溶度が大きく高温強度が高い。さらに上記範囲のγ’相の固溶と析出が可能である。そのために、Ｎｉを主体とするＮｉ基合金にする必要がある。
【００２１】
次に、上記発明を具現化するための本発明のＮｉ基超耐熱合金の各元素の作用を以下に述べる。
Ｃは、脱酸元素としての作用の他に、Ｔｉ，ＮｂおよびＴａと結び付いて、安定なＭＣ型の一次炭化物を形成し、熱間加工中にオーステナイト結晶粒の粗大化を抑制する作用と、熱間摺動性を高める作用を持つために添加する。Ｃの効果は、０．００２％の添加から発揮されるが、０．１５％を越える過度の添加は、ＭＣ型炭化物の連鎖状組織を形成し、その部分がヒートクラック発生の起点となり、工具寿命を低下させる。よってＣは、重量％で０．００２〜０．１５％の添加とする。好適には０．０１〜０．０７％である。
【００２２】
Ｓｉは脱酸剤として添加され、酸化皮膜の密着性改善にも効果がある。しかし、２％を越える過度の添加は、熱間加工性の低下と常温の延性の低下を招く。よってＳｉは、重量％で２％以下とする。好適には、０．７％以下である。
Ｍｎは、脱酸剤として添加されるが、３％を越える過度の添加は、高温強度の低下を招くため、重量％で３％以下の添加とする。好適には、１％以下である。
Ｃｒは、高温加熱中に合金の表面に密着性の高い酸化皮膜を形成し、耐酸化性を高めるとともに熱間摺動性も高めることができる。この効果のために、最低１０％を越える添加を必要とするが、２５％を越える過度の添加は、σ相の析出とそれに伴う延性の低下を招くので、Ｃｒは、重量％で１０％を越え、２５％以下の範囲とする。好適には、１３〜２０％の範囲である。
【００２３】
Ｗは、α−Ｗの析出と高温強度の高いオーステナイト相を得るために、必須の添加元素である。これらの効果を得るために、Ｗは、重量％で最低１０％の添加を必要とするが、３０％を越える過度の添加は、α−Ｗの過度の析出と耐酸化性、酸化皮膜の密着性の低下を招くので、Ｗは、１０〜３０％の範囲とする。好適には、１２〜２２％の範囲である。
Ｍｏは、Ｗと同族の元素であり、Ｗの一部をＭｏに置換して、Ｗと同様の作用をもたらすことができる。しかし、その効果は、Ｗには及ばないので、Ｍｏは、重量％で１０％以下の範囲で、かつＷ＋２Ｍｏ≦３０の範囲とする。
Ｆｅは、本合金において、必ずしも添加する必要はないが、Ｎｉを主とするオーステナイト相に固溶して熱間加工性を改善し、また省資源化と低価格化にも役立つため、必要に応じて添加することができる。しかし、過度の添加は、オーステナイト相を軟化させ、γ’相の析出量を減らし、高温強度の低下を招くので、Ｆｅは重量％で１５％以下の添加とする。好適には、２〜１０％の範囲である。
【００２４】
Ａｌは、時効処理後に安定なγ’相を形成するのに必須の添加元素であり、重量％で最低０．４％以上の添加を必要する。しかし、２．５％を越える過度の添加は、γ’相の増量を招き、熱間加工性を低下させるので、Ａｌは、０．４〜２．５％の範囲とする。好適には、０．７〜１．５％の範囲である。
Ｔｉの一部は、Ｃと結び付いて安定なＭＣ型の一次炭化物を形成し、熱間加工中にオーステナイト結晶粒の粗大化を抑制する作用と、熱間摺動性を高める。また、残部のＴｉはγ’相に固溶し、γ’相を固溶強化して高温強度向上に役立つ。そのために、Ｔｉは、重量％で最低０．４％以上の添加を必要とするが、３．５％を越える過度の添加は、熱間加工性を低下させるのと同時にγ’相を不安定化して、高温長時間使用後の強度の低下を招くので、Ｔｉは０．４〜３．５％の範囲とする。好適には、０．７〜２．５％の範囲である。
また、Ａｌ，Ｔｉは耐酸化性に対しても重要な作用を有する。すなわち、Ａｌ，Ｔｉは、表面スケール形成において、前述の粒状α−Ｗ層の内部に酸化層を形成することで、粒状α−Ｗ層の効果と相まって酸化皮膜の密着性を改善し、高温耐酸化性と熱間摺動性を改善する。
【００２５】
ＮｂとＴａは、Ｔｉと同様、一部がＣと結び付いて安定なＭＣ型の一次炭化物を形成し、熱間加工中にオーステナイト結晶粒の粗大化を抑制する作用と、熱間摺動性を高める。また、残部のＮｂとＴａはγ’相に固溶し、γ’相を固溶強化して高温強度向上に役立つので、必要に応じて添加できる。しかし、両者とも重量％で３％を越える過度の添加は、熱間加工性を低下させるので、ＮｂとＴａの添加範囲は、それぞれ３％以下とする。
ＺｒとＢは、本発明において粒界強化作用により高温の強度と延性を高めるのに有効であり、本発明合金に１種または２種を適量添加できる。その効果は少量の添加量から始まるが、ＺｒおよびＢがそれぞれ重量％で０．２％および０．０２％を越えると加熱時の初期溶融温度が低下して熱間加工性が劣化するので、ＺｒおよびＢの上限は、それぞれ０．２％および０．０２％とする。
【００２６】
ＭｇとＣａは、強力な脱酸・脱硫元素として合金の清浄度を高めるとともに、高温引張やクリープ変形時さらに熱間加工時の延性改善に役立つため、１種または２種を適量添加できる。その効果は少量の添加量から始まるが、Ｍｇ，Ｃａがそれぞれ重量％で０．０２％を越えると加熱時の初期溶融温度が低下して熱間加工性が劣化するので、ＭｇおよびＣａの上限は、それぞれ０．０２％とする。
Ｎｉは、安定なオーステナイト相を形成し、ガンマプライム相の固溶および析出の基地となる。また、Ｎｉは、多量のＷを固溶できるため、高温強度の高いオーステナイト基地が得られるため、残部とする。
【００２７】
上記元素の他、重量％で１０％以下のＣｏを本発明合金に添加してもよい。
Ｃｏは、基地のオーステナイト中に固溶し、若干の固溶強化作用を有するとともに、酸化皮膜の密着性を改善させる効果も有する。Ｃｏは、Ｎｉ基地中に固溶するのでγ’相の析出にはほとんど効果を及ぼさないので好都合である。
しかしながら、Ｃｏは高価な元素であるので多量の添加は好ましくない。
【００２８】
上記元素の他、不可避の不純物元素は以下の範囲であれば本発明合金に含まれてもよい。
Ｐ≦０．０２％，Ｓ≦０．０２％，Ｏ≦０．０３％，Ｎ≦０．０３％
望ましい範囲は次の通りである。
Ｐ≦０．０１％，Ｓ≦０．０１％，Ｏ≦０．０１％，Ｎ≦０．０１％
一方、Ｙ，ＲＥＭ，Ｈｆ等の元素は熱間加工性を低下させるので本発明合金に特に添加しなくてもよいが、酸化皮膜の密着性と耐酸化性を改善させる効果をもつので、以下の範囲で添加してもよい。
Ｙ≦０．２，ＲＥＭ≦０．２，Ｈｆ≦０．２
【００２９】
また、Ｖは高温強度を向上させる効果が本発明合金の添加元素に劣り、またＲｅは高温強度向上に寄与するが合金価格が高いことにより、本発明合金に特に添加しなくてもよいが、以下の範囲ならば添加してもよい。
Ｖ≦１％，Ｒｅ≦１％
以上述べた本発明のＮｉ基超耐熱合金は、単一の真空溶解、または真空溶解後のエレクトロスラグ再溶解や真空アーク再溶解等の精練工程を経て得られたインゴットを熱間鍛造や熱間圧延等の加工工程を通して１次製品に仕上げられる。
これらの素材は前述のようにγ’相を析出させるため、９００〜１１５０℃の固溶化処理と６００〜８５０℃の時効処理を実施したのち実用に供される。
【００３０】
【実施例】
（実施例１）
表１に示す組成の合金について、本発明合金Ｎｏ．１〜２８、比較合金Ｎｏ．４１、４２および従来合金Ｎｏ．５１〜５４について真空誘導溶解によって１５ｋｇのインゴットを溶製した後、熱間加工によって３０ｍｍ角の棒材を作成した。本発明合金Ｎｏ．１〜Ｎｏ．２８と比較合金Ｎｏ．４１、４２および従来合金Ｎｏ．５１については、原子％で表されるＣ量の２倍の値である計算ＭＣ型炭化物量と画像解析により実測したα−Ｗ量を、本発明合金Ｎｏ．１〜Ｎｏ．２８と比較合金Ｎｏ．４１、４２については、次式で表される計算γ’量を表１に併せ示す。
計算γ’＝４（０．０２６［Ｃｒ］＋０．１３［Ｍｏ］＋０．１３［Ｗ］＋０．６１［Ａｌ］＋０．６８［Ｔｉ］＋０．５［Ｎｂ］＋０．５［Ｔａ］−［Ｃ］）
（［］内の元素は原子％）
【００３１】
【表１】

【００３２】
なお、α−Ｗ量の測定は、以下に示す熱処理後の試料断面を鏡面研摩後、王水にて腐食した試料を用いた。直径５μｍ以上の粒子径のものは、光学顕微鏡により観察し、８０００ｍｍ^２の面積について画像解析を行なって、面積率を求め、また、直径５μｍ未満の粒子径のものは、走査型電子顕微鏡により観察し、８０００μｍ^２の面積について画像解析を行なって面積率を求め、それぞれ算出した両者の面積率の和をα−Ｗ量とした。
ここで、比較合金Ｎｏ．４１は本発明合金に比べてＣ量が高く、比較合金Ｎｏ．４２は計算γ’量とＣｒ量が本発明合金より低い組成である。従来合金Ｎｏ．５１は特開平３−６１３４５号に記載のＮｉ−Ｗ合金、従来合金Ｎｏ．５２はＦｅ基超耐熱合金Ａ２８６、従来合金Ｎｏ．５３は析出硬化型熱間工具鋼、従来合金Ｎｏ．５４はＪＩＳＳＫＤ６１である。
【００３３】
これらの熱間加工材のうち、本発明合金Ｎｏ．１〜２８と比較合金Ｎｏ．４１、４２については、１０５０℃×３０分保持後油冷の固溶化処理と７２０℃×８時間保持後徐冷した後６２０℃×８時間保持後空冷の時効処理を行った。また、従来合金Ｎｏ．５１は９５０℃×３０分保持後空冷の固溶化処理を行った。従来合金Ｎｏ．５２は９８０℃×３０分保持後空冷の固溶化処理と７３０℃×１６時間保持後空冷の時効処理を行った。従来合金Ｎｏ．５３は１０００℃×３０分保持後徐冷の焼入れ後、硬さ３８２ＨＶとなるごとく焼もどしを行った。従来合金Ｎｏ．５４は１０２０℃×３０分保持後空冷の焼入れ後、硬さが４４６ＨＶとなるごとく焼もどしを行った。
【００３４】
表１の合金について室温および７００℃における硬さ測定、室温および７００℃における引張試験、室温から７００℃までの熱膨張係数測定を行った。硬さ試験は、ビッカース硬度計を用い荷重９８Ｎにて測定を行い、引張試験についてはＡＳＴＭ法に基づき平行部直径６．３５ｍｍ、標点間距離２５．４ｍｍに加工した試験片を用いて実施した。
熱間での摺動性を評価するために、熱間摺動試験を行った。この試験は、直径５ｍｍ×長さ２０ｍｍの丸棒試験片を作成し、それをボール盤のチャックに固定する。試験片端面を６００℃に加熱したＪＩＳＳＮＣＭ４３９製のブロックに１５４０ｒｐｍの回転速度で３０秒間押しつける。押しつけ荷重を大きくし試験片がブロックに焼付いた荷重をもって焼付き発生荷重とした。この焼付き発生荷重が大きいほど、高負荷における熱間摺動性がよいと評価できる。
【００３５】
また、熱間での耐摩耗性を評価するために、熱間摩耗試験を行った。この試験は熱間鍛造型の摩耗をシミュレートする簡便な試験法である。
より具体的には、本願出願人の発明として特開平５−２６０５５６号に記載されている。この試験は直径１６ｍｍの丸棒試験片の端面に加熱冷却の熱サイクルを与えることと約８００℃に加熱されたＪＩＳＳ４５Ｃ製ピンを摩擦摺動させることを繰返して、試験片端面にヒートクラックを伴った摩耗を起こさせる。試験片の加熱温度は６００℃とし、ピンとの摺動は面圧１４Ｎ／ｍｍ^２を４００ｒｐｍの回転速度で５秒間、冷却は３２℃の水で３．５秒間行った。加熱、摩擦摺動、冷却を１サイクルとして２０００サイクル繰返した後、試験片端面の摩耗深さを表面粗さ計にて測定して摩耗量とし、また試験片摩耗部の断面ミクロ組織を観察して発生したヒートクラックの本数および最大長を測定した。
【００３６】
さらに、耐酸化性を評価するために、表１に示す合金より、直径８ｍｍ×長さ１５ｍｍの丸棒試験片を作成し、大気中にて９００℃×１６時間保持後室温まで空冷の処理を５回繰返した後の酸化重量変化を測定した。
上記の各種試験結果を表２にそれぞれ示す。本発明合金は７００℃の硬さ、７００℃の引張強さが比較合金や従来合金より高く高温域での機械的性質に優れることが確認された。熱膨張係数は従来合金Ｎｏ．５４に代表される熱間工具鋼と近い値を示した。従来合金Ｎｏ．５２は従来合金Ｎｏ．５４に代表される熱間工具鋼に比べて熱膨張係数が高いことがわかる。
大気中にて９００℃×１６時間保持を５回繰返した後の酸化重量変化は、従来合金Ｎｏ．５３，５４を除き、他はいずれも増量値を示し、酸化皮膜の密着性は良好である。
【００３７】
【表２】

【００３８】
ただし、比較合金Ｎｏ．４２のようにＣｒ量が低いと増量値が大きくなり、従来合金Ｎｏ．５１のようにＣｒがさらに低くなると本発明合金に比べて明らかに増量値が大きくなり、耐酸化性の向上には、本発明合金並みのＣｒ量が必要であることがわかる。一方、従来合金Ｎｏ．５３，５４のような熱間工具鋼では減量値を示した。これは酸化試験中に酸化皮膜が剥離し、酸化重量はこの剥離した酸化皮膜を除外して計算したためで、酸化皮膜の密着性が本発明合金に比べて劣ることがわかる。
【００３９】
なお、図１に本発明合金Ｎｏ．１０の酸化試験後の断面表層部の走査型電子顕微鏡によるミクロ組織ならびにミクロ組織を解説するために、観察に基づく構成相の模式図を示す。酸化皮膜を構成する相の同定は、微小部Ｘ線回折ならびにエネルギー分散型Ｘ線分析により実施した。本発明合金は、いずれも７００℃〜１０００℃の高温酸化雰囲気では、このような表面側からＣｒを主とする表面酸化層、粒状α−Ｗ層およびＡｌ，Ｔｉを主とする内部酸化層の３層構造をとる。すなわち、粒状のα−Ｗ層がＣｒを主とする表面酸化層とＡｌ，Ｔｉを主とする内部酸化層にはさまれた形になっており、このような酸化皮膜構造が、本発明合金の高い耐酸化性と酸化皮膜密着性およびそれらに付随する高い熱間摺動性と熱間での耐摩耗性を有する理由である。
【００４０】
熱間摺動試験では本発明合金は従来合金Ｎｏ．５４に代表される熱間工具鋼に比べて焼付き発生荷重が高いことが確認された。熱間摺動性は密着性および安定性に優れた酸化皮膜の生成と高温耐力が要求されるが、本発明合金はこれを満たし、熱間摺動性に優れることが確認された。
熱間摩耗試験では、本発明合金は摩耗量が従来合金Ｎｏ．５４に代表される熱間工具鋼に比べて非常に少ない結果となった。Ｗを多量に含有する従来合金Ｎｏ．５１は熱間摺動試験の焼付き発生荷重は高いが、α−Ｗ量が体積％で３０％を越えており、またＣｒ量も低いので、生成される酸化皮膜が長時間における密着性および安定性に欠けるため本発明合金に比べて耐摩耗性に劣る。従来合金Ｎｏ．５３を除いて、ヒートクラックが発生した。従来合金Ｎｏ．５３はＡＣ_１点が低いため摩擦摺動によって試験片端面の温度は容易にこの温度を越えオーステナイト化し、それに伴い摩擦摺動部の強度が極端に低下し、ヒートクラックが発生してもすぐ摩耗してしまい試験後には観察されない。
【００４１】
従来合金Ｎｏ．５２は熱膨張係数が高いためヒートクラックが発生しやすく、従ってヒートクラックが摩擦摺動面の摩擦係数を増加し摩耗が本発明合金より多くなった。Ｃ量の高い比較合金Ｎｏ．４１はＭＣ型炭化物の連鎖状組織を形成しており、その部分がヒートクラックの起点となっており、ヒートクラックが多く発生し、それに伴い摩耗も多い。比較合金Ｎｏ．４２はγ’量が原子％で１０％より低いため７００℃の引張強さが本発明合金より低く摩耗が多い。本発明合金はこの試験によりヒートクラックを伴った熱間での耐摩耗性に優れることが確認された。
【００４２】
（実施例２）
表１に示す本発明合金Ｎｏ．４と従来合金Ｎｏ．５２，５３について実施例１と同じ熱処理を行い、ボールエンドミルを用いた被削性試験を行った。マシニングセンターを用い、セラミックコーティング超硬製ボールエンドミルにより、水溶性切削油を使用し、刃先摩耗量の測定により、被削性を調査した。試験条件は、切削速度５０ｍ／ｍｉｎ、送り０．０５ｍｍ／刃、切り込み量は、軸方向２ｍｍ、外周刃方向１ｍｍ、切削長２５ｍの条件で実施し、エンドミルの切刃逃げ面摩耗幅を測定した。試験結果を表３に示す。
本発明合金は、Ｎｉ基超耐熱合金の範疇に属するにもかかわらず、被削性が従来のＦｅ基超耐熱合金Ｎｏ．５２はいうまでもなく、従来の熱間工具鋼Ｎｏ．５３に比べても優れた被削性を示した。この良好な被削性は微細α−Ｗ相の分散によってもたらされる。このように本発明合金は、時効処理後でも十分な被削性を有するので、金型加工後の熱処理が不要なプリハードン工具としても使用できる。
【００４３】
【表３】

【００４４】
（実施例３）
以下、本発明合金を熱間鍛造金型に実施した例を示す。
表１に示す、本発明合金Ｎｏ．１０、Ｎｏ．１４の類似成分で溶製し直したもの、および従来合金Ｎｏ．５４の組成のものの素材を準備し、これから熱間鍛造金型を製作し、実用テストを行った結果を表４に示す。金型は自動車部品のギヤシャフトを製作する型であり、寸法は直径８０ｍｍ×長さ１６０ｍｍである。このギヤシャフトを製作する金型は粗地成形用の粗地型と仕上成形用の仕上型よりなるが、摩耗の激しい粗地型にてテストを行った。ワーク材はＳ３５Ｃであり、高周波加熱装置により１２００℃に加熱する。鍛造成形は最大能力１６００ｔのクランクプレスを使用し、各型打ち後に白色系潤滑剤を型表面に噴霧した。
従来合金Ｎｏ．５４はＪＩＳＳＫＤ６１である。熱処理は、熱間鍛造金型に荒加工した後、１０２０℃に加熱し、２００℃の油に浸漬する油焼入れを行い、硬さが４４６ＨＶとなるごとく焼もどしを行った。その後仕上加工を行い、実用テストに供した。
【００４５】
【表４】

【００４６】
表４にこれら合金で製作した熱間鍛造金型の寿命を示す。この熱間鍛造作業により金型は高温のワークとの接触およびワークとの摺動により金型表層部は熱影響を大きく受けるが、本金型は従来、黒鉛系潤滑剤による作業が、白色系潤滑剤に変更になったことに伴い、特にワークとの摺動発熱作用が大きくなったため、従来合金Ｎｏ．５４は熱影響により焼もどしマルテンサイトの軟化およびＡＣ_１変態点を越えることによりオーステナイト変態を起こし、表層部の強度が低下し摩耗が起こりやすく早期に寿命となった。また白色系潤滑剤が各型打ち後に型表面に噴霧されるため、型表面には加熱冷却の熱サイクルが負荷され、型表面にヒートクラックが発生する。金型表面に発生したヒートクラックはワークとの摩擦係数を増加するため、ワークとの摺動発熱が大きくなり型表層部の熱影響を増加させ、また金型表層部に発生する剪断応力が大きくなり型表層部の塑性流動が増加するため、摩耗を促進させる。従来合金Ｎｏ．５４はヒートクラック発生により熱影響が大きくなり、発生する剪断応力が型表層部強度を越えて摩耗が起こりやすく早期に寿命となった。
【００４７】
本発明合金Ｎｏ．１０、Ｎｏ．１４は、その優れた酸化被膜特性によりワークとの摺動発熱を抑制するとともに、熱影響を受けた場合でも高温の強度が鍛造作業により発生する剪断応力より勝っているので、表層部の塑性流動は起こらず、摩耗が起こりにくく、寿命が延びた。
本発明合金Ｎｏ．１０、Ｎｏ．１４にもヒートクラックは発生した。ヒートクラックはＮｏ．１４の方が多く発生しており、寿命も短い。Ｎｏ．１４はα−Ｗ量の多い合金であるが、実用テスト後の金型の摩耗部のミクロ組織を観察した結果、ヒートクラックはα−Ｗに沿って発生しており、α−Ｗ量の多い点でヒートクラックが発生しやすく摩耗が多くなった。本実用テストではヒートクラックの発生が摩耗と関係深いが、潤滑の不十分な金型での適用においてはα−Ｗの熱間摺動時の耐焼付き性の効果が表われた。
【００４８】
本発明合金は熱間での硬さでは従来合金に勝るが、室温の硬さでは従来合金に劣る場合がある。その場合、厚みのある金型では鍛造作業時にプレスの負荷を受けて金型が変形する場合があるが、型彫り部に本発明合金を、土台部分を従来合金Ｎｏ．５４やＪＩＳＳＫＴ４等を組み合わせた金型にすることにより、本発明合金の耐摩耗性向上の効果が表われた。
以上より、本発明合金は主に７００〜９００℃程度の温度域で作業することの多い熱間鍛造型として優れた性能を有することが明らかとなった。
【００４９】
（実施例４）
以下、本発明合金をＣｕ合金押出し工具に実施した例を示す。
表１に示す、本発明合金Ｎｏ．１０の類似成分で溶製し直したもの、および従来合金Ｎｏ．５２の組成のものの素材を準備し、これから熱間押出し工具を製作し、実用テストを行った結果を表５に示す。表１に示す従来合金Ｎｏ．５２はＪＩＳＳＵＨ６６０でありＡ２８６の名で知られている。この合金はＣｕまたはＣｕ合金の熱間押出し工具として知られている。
【００５０】
【表５】

【００５１】
熱間押出し工具は焼ばめによる二重構造のものを用い、外筒にＪＩＳＳＫＴ４を用い、内筒を本発明合金Ｎｏ．１０製と従来合金Ｎｏ．５２製のものを製作し比較した。外筒は外径２００ｍｍ、内筒は外径１００ｍｍ、内径６０ｍｍとし、長さはともに２００ｍｍの小型の二重構造の工具を本発明合金Ｎｏ．１０製と従来合金Ｎｏ．５２製について製作した。これらの工具を用いて１００ｔプレスにより、９５０℃の純銅ビレットの押出し実験を行った。内筒は８００℃程度の高温と５００Ｎ／ｍｍ^２前後の高圧にさらされ、熱応力により亀甲状のヒートクラックが生じ、表面剥離を起こして寿命となる。
従来合金Ｎｏ．５２の場合、約１００００ケ成形時に、既に内径面にヒートクラックの発生が認められたが、本発明合金Ｎｏ．１０の場合は、約３００００ケ成形後に僅かにヒートクラックが認められる程度であった。
本発明合金Ｎｏ．１０は従来合金Ｎｏ．５２に比べて、高温域での引張強さが大きい上、熱膨張係数が小さいのでヒートクラックが発生しにくく、このことが押出し工具の寿命を飛躍的に向上させた。この結果より、本発明合金は熱間押出し工具として優れた性能を有することが明らかとなった。
【００５２】
【発明の効果】
本発明合金は、熱間摺動性、熱間での耐摩耗性、耐ヒートクラック性や耐酸化性等の高温特性と被削性および熱間加工性に優れた合金であり、自動車用のクランクシャフトやコンロッド等自動車部品を成形する温・熱間鍛造金型や、ＣｕやＡｌまたはそれらの合金の熱間押し出し工具等に適用すれば、従来材に比較して工具の寿命を飛躍的に向上させることができる。
【図面の簡単な説明】
【図１】本発明合金Ｎｏ．１０を、大気中にて９００℃×１６時間保持後、室温まで空冷処理を５回繰り返した後の断面表層部の走査型電子顕微鏡による金属ミクロ組織写真ならびに構成相の模式図を示す図である。[0001]
[Industrial applications]
The present invention is mainly used for warm and hot forging dies used in the hot temperature range of 900 ° C. or less, hot forging dies, hot extrusion tools and parts of Cu, Al, or an alloy thereof. A Ni-based super heat-resistant alloy.
[0002]
[Prior art]
In response to the recent trend toward the common use of parts in the automobile industry and a reduction in the number of work steps, it has been desired to improve the life of hot and hot forging dies for molding automobile parts such as crankshafts and connecting rods for automobiles. However, at present, hot tool steel represented by the conventional SKD61 has not been able to meet the needs. In addition, Fe-based super heat-resistant alloys such as A286 and Inconel 718 alloy are used for hot-extrusion tools of Cu, Al or an alloy thereof, but in this application, the life is improved. Reduction of man-hours is required.
In addition to these, Japanese Patent Application Laid-Open No. 3-61345 discloses a hot tool used in a high temperature range of about 1200 ° C., such as a perforated plug and a mandrel used for manufacturing a seamless steel pipe, and a high temperature of 1000 ° C. to 1150 ° C. Japanese Patent Application Laid-Open No. 4-41641 discloses a Ni-based super heat-resistant alloy for a constant temperature forging die used in an air atmosphere.
Further, although the purpose is different from the present invention, examples of super heat-resistant alloys based on Ni-Cr-W include, for example, JP-B-54-33212, JP-B-60-58773, JP-B-58-502, and JP-B-562. No. 1061 is known.
[0003]
[Problems to be solved by the invention]
Forging dies used in hot or warm conditions, or extruded tools made of Cu, Al or their alloys, are capable of hot working, have excellent machinability, and have high proof stress and toughness at room temperature, and 700 Excellent hot slidability in a high temperature range of about 900C to 900C, hot wear resistance and heat crack resistance are required.
Here, in order to improve hot slidability and hot abrasion resistance, it is required to increase adhesion and stability of the oxide film and high-temperature proof stress. However, in the conventional SKD61 class hot work tool steel, when the temperature exceeds 700 ° C., the tempered martensite softens and the high-temperature proof stress decreases, and at the same time, the Cr content in the steel is low, so that the adhesion and the stability are excellent. An oxide film is not formed, and the workpiece and the tool come into direct contact, causing early abrasion and finally seizure.
[0004]
On the other hand, improvement in heat crack resistance requires a reduction in the coefficient of thermal expansion and a high temperature proof stress. However, Fe-base superalloys such as A286 have a high coefficient of thermal expansion, and the gamma prime phase, which is a precipitation strengthening phase, is Ni ₃ Due to the composition mainly composed of Ti, it transforms into an eta phase which is a stable phase in a temperature range exceeding 700 ° C., and as a result, there is a problem that high-temperature strength is reduced.
Further, a hot tool is required to have excellent machinability, but a super heat-resistant alloy such as A286 generally has a disadvantage that the machinability is much worse than that of a martensitic material.
By the way, some super heat-resistant alloys for Ni-based hot tools have been proposed. For example, the alloy disclosed in Japanese Patent Application Laid-Open No. 3-61345 has insufficient oxidation resistance, and is not suitable for hot working. Has insufficient abrasion resistance. Further, the alloy disclosed in Japanese Patent Application Laid-Open No. 4-41641 has a problem that the hot slidability at a high sliding speed is insufficient because the alloy is intended for a constant temperature forging die.
[0005]
In addition, the above-mentioned Ni-Cr-W-based Ni-base superalloys are premised on using heat resistance and high creep strength at high temperatures in the extremely high temperature range of around 1000 ° C for the first purpose. And Therefore, the conventional Ni-based super heat-resistant alloys to which a large amount of W is added have a high-temperature creep strength increased by solid solution strengthening of W, and the high strength at 700 to 900 ° C., which is the object of the present invention, It is not sufficient for improvement of hot slidability, and it has been difficult to use it in characteristics intended in the field of use intended by the present invention.
[0006]
[Means for Solving the Problems]
In view of the problems of these hot tool steels and Fe-base and Ni-base super heat-resistant alloys, the present inventors can perform hot working with one of the main objects being hot tool steel replacement, and have oxidation resistance, The inventor has worked diligently on the invention of a material capable of improving high-temperature strength, coefficient of thermal expansion, and machinability. As a result, an austenitic phase mainly composed of Ni containing W up to the solid solubility limit is used as a main phase, and Ni precipitated by aging treatment ₃ By finding the optimal content of a gamma prime phase mainly composed of (Al, Ti), a solid solution of W (α-W) and an MC type primary carbide, hot slidability particularly suitable for a hot tool, The inventors have invented a Ni-based super heat-resistant alloy having excellent hot wear resistance, heat crack resistance, and oxidation resistance. Further, in the alloy having this structure, the surface scale structure generated in a high-temperature atmosphere is such that, from the surface side, an external oxide layer mainly composed of Cr, a granular α-W layer, and an internal oxide mainly composed of Al and Ti are formed. It has been newly found that a layer having a structure of at least three or more layers improves the adhesion of the oxide film and is useful for improving high-temperature oxidation resistance and hot slidability.
[0007]
The features of the Ni-base superalloy according to the present invention do not exist only in the alloy composition, as can be understood from the constituent features of the first and second inventions described below. That is, the critical difference between the Ni-base superalloy of the present invention and the conventional W-containing Ni-base superalloy is that the present invention has an age-treated structure characterized by a gamma prime phase. . Since the aging treatment is performed at 600 to 850 ° C, in applications where the temperature is around 1000 ° C or higher, such as conventional Ni-based super heat-resistant alloys, the aging treatment phase beforehand is meaningless because the precipitating age-hardened phase forms a solid solution. And was not recognized.
In the present invention, the age-precipitated phase, which could not be adopted with the conventional Ni-base superalloy having a similar composition, is positively utilized. Therefore, the structure of the Ni-base superalloy of the present invention is characterized by the presence of a large amount of gamma prime phase of 10 to 30%.
[0008]
That is, Ming is Having a structure that has been subjected to aging treatment, and the phase constituting the structure is 0.02 to 1.5% in terms of the amount of MC type carbide calculated from the atomic% according to the formula (1), and from the atomic% according to the formula (2). A Ni-base comprising a calculated gamma prime amount of 10 to 30%, and an α-W amount of 0.5 to 30% expressed by volume%, with the balance being an austenite phase mainly composed of Ni. It is a super heat resistant alloy.
(1) 2 [C]
(2) 4 (0.026 [Cr] +0.13 [Mo] +0.13 [W] +0.61 [Al] +0.68 [Ti] +0.5 [Nb] +0.5 [Ta]-[C ])
(However, in the above formula, the element without addition is calculated as 0)
[0009]
The chemical composition of the Ni-base superalloy for embodying the present invention is, as a percentage by weight, C 0.002 to 0.15%, Si 2% or less, Mn 3% or less, Cr more than 10% and 25% or less, W10 to W10. It contains 30%, Fe 15% or less, Al 0.4-2.5%, Ti 0.4-3.5%, and the balance consists of unavoidable impurities and Ni.
[0010]
Up The composition of the desired alloy of the present invention described above can further substitute one or two of Ni and 3.0% or less of Nb and 3.0% or less by weight by substituting a part of Ni. Further, in addition to the above composition, a part of Ni is substituted, and 10% or less by weight of Mo is added in the range of W + 2Mo ≦ 30, and Zr of 0.2% or less by weight is further added. And 0.02% or less of B or one or two kinds of B, and 0.02% or less of Mg and 0.02% or less of one or two kinds of Ca by weight% can be added. . Of the above alloy elements, those described below Nb can be selectively added in all possible combinations. In addition, the selectively added element must satisfy the above-described equation for estimating the amount of gamma prime.
[0011]
[Action]
The function of the structure defined by the Ni-base superalloy of the present invention will be described.
The MC type primary carbide is the hardest phase in the Ni-base superalloy of the present invention, and functions to suppress austenite crystal grains from being coarsened during hot working in the production of the present alloy, Has the effect of enhancing slidability. Assuming that the total amount of added C is present in association with Ti, Nb and Ta, twice the amount of C expressed in atomic% can be regarded as the content of MC type carbide. it can. Therefore, the above-mentioned amount of MC type carbide is a value calculated by 2 [C] in the equation (1). In order to exhibit the effect of the MC type carbide, it is necessary to add the amount of MC exceeding at least 0.02% in atomic%, but excessive addition exceeding 1.5% may cause the chain structure of the MC type carbide. Is formed, and that portion serves as a starting point of the occurrence of a heat crack, thereby shortening the tool life. Therefore, the amount of MC-type carbide calculated from atomic% is 0.02 to 1.5%. The preferred amount of MC carbide is 0.1-1.0%.
[0012]
Precipitation of the gamma prime (γ ′) phase by aging has the greatest effect on improving the strength at room temperature and high temperature. The basic composition of the gamma prime phase is Ni ₃ Although represented by Al, in a multi-component alloy, various additional elements are actually dissolved in Al sites in the gamma prime phase. Since it is very difficult to measure the amount of precipitation of the gamma prime phase, in the present invention, literature values (Hirata Harada, Michio Yamazaki: “γ ′ precipitation strengthened Ni-based heat-resistant casting containing Ti, Ta, W”) Alloy Design ", Iron and Steel, vol. 65, (1979), p1059), and the alloy No. 1 of the present invention described in Example 1 described later. The composition of γ phase and γ ′ phase and the ratio of both phases were determined, and the coefficient of each element entering the Al site of γ ′ phase was determined from atomic%. Here, from the component range of the alloy of the present invention, Ni ₃ Cr, Mo, W, Ti, Nb, and Ta were considered as elements that form a solid solution in the Al site of Al. However, in the case of No. 1 of the first embodiment. 5 is a component system that does not include Mo and Ta, so that the coefficients of Mo and Ta cannot be calculated. Therefore, the coefficients of W and Nb of the same family are used as the coefficients occupying the Al site of the γ ′ phase of Mo and Ta, respectively.
[0013]
In the calculation, Ti, Nb, and Ta were assumed to be those in which the total amount of added C and the remainder forming the MC-type carbide participated in the equilibrium of the γ-γ ′ phase. From the calculation results, the total amount of the added elements that enter the Al site of the γ ′ phase was determined, and four times this was defined as the calculated γ ′ amount. That is, the gamma prime amount in the present invention is calculated by the following equation, and corresponds to the above-mentioned equation (2).
Calculation γ ′ = 4 (0.026 [Cr] +0.13 [Mo] +0.13 [W] +0.61 [Al] +0.68 [Ti] +0.5 [Nb] +0.5 [Ta] − [ C])
(Elements in [] indicate atomic%, and elements without addition are calculated as 0.)
[0014]
Next, in the Ni-base superalloy of the present invention, a heat treatment for precipitating the γ 'phase is performed. In this heat treatment, first, a solution treatment is performed to make the structure after hot working uniform and to cause a precipitate to form a solid solution once. The temperature of the solid solution strengthening is desirably in the range of 900 to 1200 ° C. At 900 ° C. or lower, the solid solution of the precipitate is insufficient, and at 1200 ° C. or higher, the crystal grains become coarse, which is disadvantageous in strength and heat crack resistance. However, when emphasis is placed on the strength at room temperature, the solution treatment may be performed at a lower temperature or may be omitted in order to leave distortion of forging.
[0015]
The aging treatment is subsequently performed, and this aging treatment is indispensable for the Ni-based super heat-resistant alloy of the present invention to precipitate the γ 'phase and increase the high-temperature strength. The aging treatment is preferably performed at a temperature of 600 to 850C. If the temperature is lower than 600 ° C., a large amount of time is required for the precipitation of the γ ′ phase, and if the temperature exceeds 850 ° C., the γ ′ phase to be deposited becomes coarse or the γ ′ phase does not precipitate while being in a solid solution. High temperature strength cannot be obtained.
This aging treatment can be carried out once or twice or more, but a total treatment time of at least 3 Hr is required.
When the calculated γ ′ amount precipitated by the aging treatment is less than 10%, the high-temperature strength is reduced, and accordingly, the resistance to hot wear and hot sliding is reduced, and the tool life is reduced. Conversely, if the calculated γ ′ amount exceeds 30%, the hot deformation resistance increases, and hot forging becomes difficult. Therefore, the amount of γ ′ calculated by the above equation is in the range of 10 to 30%. Preferably, it is in the range of 13 to 25%.
[0016]
The alloys described in JP-B-54-33212, JP-B-60-58773, JP-B-58-502, and JP-B-56-21061 are intended to strengthen solid solution by adding W, and gamma prime. Almost no forming elements. Alternatively, even when a small amount is contained, the aging treatment is not performed, so that no enhancement by the gamma prime phase is attempted. This is because the use temperature is around 1000 ° C., so that the γ ′ phase forms a solid solution during use. In this case, the action of the γ ′ forming elements Al, Ti, Nb, etc. Oxidation resistance is improved, and Ti and Nb are not intended to positively strengthen by the γ 'phase as in the present invention.
Furthermore, since each of the above alloys emphasizes high-temperature creep strength near 1000 ° C., it is necessary to carry out the solution treatment at a temperature of 1200 ° C. or more, or to perform a re-heating at 1080 ° C. or more after the solution treatment. Since it is used, the crystal grains are coarsened, which is disadvantageous for applications requiring strength at 700 to 900 ° C. and heat crack resistance, which are the objects of the present invention.
[0017]
Next, the uniformly dispersed α-W phase (solid solution of W having a bcc structure) is an important alloy phase essential for the alloy of the present invention. The α-W phase contributes to a reduction in the coefficient of thermal expansion of the alloy and also prevents seizure during hot sliding. Further, as described below, α-W also has an effect of improving the adhesion of the oxide film. The effect of α-W is exhibited from 0.5% or more in terms of volume% on the real side such as image analysis, but if it exceeds 30%, the oxidation resistance and the hot workability are reduced. Therefore, the content of α-W is in the range of 0.5 to 30% by volume. Preferably, it is in the range of 3-20%.
A feature of the Ni-based super heat-resistant alloy of the present invention is that a γ 'phase and an α-W phase are present at the same time, which is not found in a conventional super heat-resistant alloy.
That is, a normal γ ′ precipitation-strengthened superalloy naturally has a γ ′ phase but no α-W phase.
[0018]
Among conventional Ni-base superalloys, alloys with a high W may have an α-W phase, but at most, they precipitate a small amount at crystal grain boundaries and in grains to improve creep strength. This is completely different from the operation and effect of the present invention. Further, these conventional alloys have a decisive difference in that the γ 'phase does not exist as described above. In the alloy of the present invention, the required properties such as strength at 700 to 900 ° C., hot sliding, hot abrasion, heat crack resistance, oxidation resistance and the like are first obtained by the simultaneous existence of the γ ′ phase and the α-W phase. At the same time, it became possible to be satisfied.
Further, in the Ni-base superalloy of the present invention, for example, as shown in FIG. 1 in Example 1, the surface scale structure generated in a high-temperature atmosphere has an external oxide layer mainly composed of Cr, a granular α It is characterized in that it has a structure of at least three layers of a -W layer and an internal oxide layer mainly composed of Al and Ti. The mechanism of this oxide film formation is not clear, but if the oxide film structure does not take at least three layers, the adhesion and oxidation resistance of the oxide film will be reduced.
[0019]
For example, in the case of the Ni-base superalloy of the present invention, if α-W particles exceeding 5 μm are exposed on the surface, the portion becomes WO ₃ Oxide film is formed, and the portion does not have sufficient adhesion, but the total amount of α-W is suppressed to 30% or less as a whole, so that the life of the tool itself is greatly reduced. Not be. The granular α-W layer in the scale is changed from the α-W phase of the precipitated phase, and it is desirable to increase the particle diameter to 1 μm or less. Therefore, the structure of the surface scale when the alloy of the present invention is heated in the air has at least the following components from the surface side: an external oxide layer mainly composed of Cr, a granular α-W layer, and an internal oxide layer mainly composed of Al and Ti. It is desirable to have a structure of three or more layers.
[0020]
Further, except for the above amount of MC type carbide, the γ 'phase and the α-W phase, and the unavoidable impurities, the balance is an austenite phase mainly composed of Ni. This matrix has a high solid solubility of W and a high high-temperature strength. Further, solid solution and precipitation of the γ 'phase in the above range are possible. Therefore, it is necessary to use a Ni-based alloy mainly composed of Ni.
[0021]
Next, the action of each element of the Ni-base superalloy of the present invention for realizing the above-described invention will be described below.
C combines with Ti, Nb and Ta in addition to the action as a deoxidizing element to form a stable MC type primary carbide, and to suppress the coarsening of austenite crystal grains during hot working; It is added because it has the effect of enhancing hot slidability. The effect of C is exhibited from the addition of 0.002%, but excessive addition exceeding 0.15% forms a chain structure of MC type carbide, and that portion becomes a starting point of heat crack generation, and the tool becomes a starting point. Reduce life. Therefore, C is added in an amount of 0.002 to 0.15% by weight. Preferably it is 0.01 to 0.07%.
[0022]
Si is added as a deoxidizing agent, and is also effective in improving the adhesion of the oxide film. However, excessive addition exceeding 2% causes a reduction in hot workability and a reduction in ductility at room temperature. Therefore, Si is set to 2% or less by weight%. Preferably, it is 0.7% or less.
Mn is added as a deoxidizing agent, but excessive addition exceeding 3% causes a decrease in high-temperature strength. Preferably, it is 1% or less.
Cr forms an oxide film having high adhesion on the surface of the alloy during heating at a high temperature, thereby improving oxidation resistance and hot slidability. For this effect, addition of at least 10% is required, but excessive addition of more than 25% causes precipitation of the σ phase and accompanying decrease in ductility. Exceed and within the range of 25% or less. Preferably, it is in the range of 13-20%.
[0023]
W is an essential additive element in order to precipitate α-W and obtain an austenite phase having high high-temperature strength. In order to obtain these effects, W needs to be added at least 10% by weight, but excessive addition exceeding 30% causes excessive precipitation of α-W, oxidation resistance, and adhesion of oxide film. Therefore, W is set to be in the range of 10 to 30%, since this causes a reduction in the properties. Preferably, it is in the range of 12 to 22%.
Mo is an element of the same family as W, and can substitute a part of W with Mo to provide the same action as W. However, since the effect is not as great as W, Mo is limited to a range of 10% or less by weight and a range of W + 2Mo ≦ 30.
Fe is not always necessary to be added to the present alloy, but it is necessary because it forms a solid solution with the austenite phase mainly containing Ni to improve hot workability, and also contributes to resource saving and cost reduction. It can be added accordingly. However, excessive addition softens the austenite phase, reduces the precipitation amount of the γ 'phase, and lowers the high-temperature strength. Therefore, Fe is added in an amount of 15% by weight or less. Preferably, it is in the range of 2 to 10%.
[0024]
Al is an essential element for forming a stable γ 'phase after aging treatment, and requires at least 0.4% or more by weight of Al. However, excessive addition exceeding 2.5% causes an increase in the amount of the γ 'phase and lowers hot workability. Therefore, Al is set in the range of 0.4 to 2.5%. Preferably, it is in the range of 0.7-1.5%.
Part of Ti combines with C to form stable MC-type primary carbides, thereby suppressing the coarsening of austenite crystal grains during hot working and enhancing hot slidability. Further, the remaining Ti forms a solid solution in the γ ′ phase, and solid solution strengthens the γ ′ phase to help improve the high-temperature strength. Therefore, Ti needs to be added in an amount of at least 0.4% by weight or more, but excessive addition exceeding 3.5% reduces the hot workability and simultaneously destabilizes the γ ′ phase. Therefore, Ti is reduced to a range of 0.4 to 3.5% because of the decrease in strength after long-term use at high temperatures. Preferably, it is in the range of 0.7-2.5%.
Al and Ti also have an important effect on oxidation resistance. That is, Al and Ti form an oxide layer inside the above-mentioned granular α-W layer in the formation of the surface scale, thereby improving the adhesion of the oxide film in combination with the effect of the granular α-W layer, and And hot sliding properties are improved.
[0025]
Nb and Ta, like Ti, partially combine with C to form a stable MC-type primary carbide, and have an effect of suppressing coarsening of austenite crystal grains during hot working and a hot sliding property. Enhance. Further, the remaining Nb and Ta form a solid solution in the γ ′ phase, and solid solution strengthening of the γ ′ phase is useful for improving the high-temperature strength, and thus can be added as necessary. However, in both cases, excessive addition exceeding 3% by weight reduces the hot workability. Therefore, the addition ranges of Nb and Ta are set to 3% or less, respectively.
Zr and B are effective for enhancing the high-temperature strength and ductility by the grain boundary strengthening action in the present invention, and one or two or more of them can be added to the alloy of the present invention in an appropriate amount. The effect starts with a small amount of addition. However, if the content of Zr and B exceeds 0.2% and 0.02% by weight, respectively, the initial melting temperature at the time of heating lowers and the hot workability deteriorates. The upper limits of Zr and B are set to 0.2% and 0.02%, respectively.
[0026]
Mg and Ca serve as powerful deoxidizing and desulfurizing elements to enhance the cleanliness of the alloy and to help improve ductility during high-temperature tensile and creep deformation and hot working, so that one or two of them can be added in an appropriate amount. The effect starts with a small amount of addition. However, if the content of Mg and Ca exceeds 0.02% by weight, respectively, the initial melting temperature at the time of heating decreases and the hot workability deteriorates. Is 0.02% each.
Ni forms a stable austenite phase and serves as a base for solid solution and precipitation of the gamma prime phase. In addition, Ni can form a solid solution of a large amount of W, so that an austenite matrix having high high-temperature strength can be obtained.
[0027]
In addition to the above elements, 10% or less by weight of Co may be added to the alloy of the present invention.
Co forms a solid solution in the austenite of the matrix, has a slight solid solution strengthening action, and also has an effect of improving the adhesion of the oxide film. Co is advantageous because it forms a solid solution in the Ni matrix and has little effect on the precipitation of the γ 'phase.
However, since Co is an expensive element, it is not preferable to add a large amount of Co.
[0028]
In addition to the above elements, unavoidable impurity elements may be included in the alloy of the present invention as long as the following ranges are satisfied.
P ≦ 0.02%, S ≦ 0.02%, O ≦ 0.03%, N ≦ 0.03%
Desirable ranges are as follows.
P ≦ 0.01%, S ≦ 0.01%, O ≦ 0.01%, N ≦ 0.01%
On the other hand, elements such as Y, REM, and Hf reduce the hot workability and do not need to be particularly added to the alloy of the present invention, but have the effect of improving the adhesion and oxidation resistance of the oxide film. May be added.
Y ≦ 0.2, REM ≦ 0.2, Hf ≦ 0.2
[0029]
Further, V is inferior to the additive element of the alloy of the present invention in improving the high-temperature strength, and Re contributes to the improvement of the high-temperature strength but is not necessarily added to the alloy of the present invention because of the high alloy price. You may add in the following ranges.
V ≦ 1%, Re ≦ 1%
The Ni-based super heat-resistant alloy of the present invention described above can be used for hot forging or hot ingot obtained through a single vacuum melting or a refining process such as electroslag remelting or vacuum arc remelting after vacuum melting. It is finished into a primary product through processing steps such as rolling.
These materials are put to practical use after a solution treatment at 900 to 1150 ° C and an aging treatment at 600 to 850 ° C to precipitate the γ 'phase as described above.
[0030]
【Example】
(Example 1)
With respect to the alloys having the compositions shown in Table 1, the alloy No. of the present invention was used. Nos. 1 to 28, Comparative alloy Nos. Nos. 41 and 42 and conventional alloy Nos. After producing 15 kg of ingots for 51 to 54 by vacuum induction melting, rods of 30 mm square were prepared by hot working. The alloy No. of the present invention. 1 to No. No. 28 and Comparative Alloy No. 28 Nos. 41 and 42 and conventional alloy Nos. For alloy No. 51 of the present invention, the calculated MC-type carbide amount, which is twice the C amount expressed in atomic%, and the α-W amount actually measured by image analysis were used. 1 to No. No. 28 and Comparative Alloy No. 28 For 41 and 42, the calculated γ ′ amount represented by the following equation is also shown in Table 1.
Calculation γ ′ = 4 (0.026 [Cr] +0.13 [Mo] +0.13 [W] +0.61 [Al] +0.68 [Ti] +0.5 [Nb] +0.5 [Ta] − [ C])
(Elements in [] are atomic%)
[0031]
[Table 1]

[0032]
In the measurement of the amount of α-W, a sample obtained by subjecting the cross section of the sample after the heat treatment described below to mirror polishing and corroding with aqua regia was used. Particles having a diameter of 5 μm or more were observed with an optical microscope, and 8000 mm ² The area ratio was determined by performing image analysis on the area of, and those having a particle diameter of less than 5 μm were observed with a scanning electron microscope to obtain 8000 μm. ² The area ratio was determined by performing image analysis on the area of, and the sum of the calculated area ratios was defined as the α-W amount.
Here, the comparative alloy No. No. 41 has a higher C content than the alloy of the present invention, and Comparative Alloy No. 41 Reference numeral 42 denotes a composition in which the calculated γ ′ amount and the Cr amount are lower than those of the alloy of the present invention. Conventional alloy No. No. 51 is a Ni-W alloy described in JP-A-3-61345, and a conventional alloy No. 52 is Fe-based super heat-resistant alloy A286, conventional alloy No. 53 is a precipitation hardening type hot work tool steel, conventional alloy No. 53; 54 is JIS SKD61.
[0033]
Among these hot worked materials, the alloy No. 1 of the present invention was used. Nos. 1 to 28 and Comparative Alloy Nos. As for 41 and 42, a solution treatment of oil cooling after holding at 1050 ° C. for 30 minutes, a slow cooling after holding at 720 ° C. for 8 hours, an aging treatment of air cooling after holding at 620 ° C. for 8 hours were performed. In addition, the conventional alloy No. No. 51 was subjected to a solution treatment of air cooling after holding at 950 ° C. × 30 minutes. Conventional alloy No. Sample No. 52 was subjected to a solution treatment of air cooling after holding at 980 ° C. × 30 minutes and an aging treatment of air cooling after holding at 730 ° C. × 16 hours. Conventional alloy No. No. 53 was kept at 1000 ° C. for 30 minutes, then quenched by slow cooling, and then tempered to a hardness of 382 HV. Conventional alloy No. No. 54 was kept at 1020 ° C. for 30 minutes, air-quenched, and then tempered to a hardness of 446 HV.
[0034]
The hardness of the alloys in Table 1 at room temperature and 700 ° C., the tensile test at room temperature and 700 ° C., and the coefficient of thermal expansion from room temperature to 700 ° C. were measured. The hardness test was performed at a load of 98 N using a Vickers hardness tester, and the tensile test was performed using a test piece processed to a parallel part diameter of 6.35 mm and a gauge length of 25.4 mm based on the ASTM method. .
A hot sliding test was performed to evaluate hot sliding properties. In this test, a round bar test piece having a diameter of 5 mm and a length of 20 mm is prepared and fixed to a chuck of a drilling machine. The test piece end face is pressed against a block made of JIS SNCM439 heated to 600 ° C. at a rotation speed of 1540 rpm for 30 seconds. The pressing load was increased, and the load at which the test piece seized the block was defined as the seizure load. It can be evaluated that the larger the seizure load, the better the hot slidability under a high load.
[0035]
A hot wear test was performed to evaluate hot wear resistance. This test is a simple test that simulates the wear of a hot forging die.
More specifically, the invention of the present applicant is described in JP-A-5-260556. In this test, a heat cycle of heating and cooling was given to the end face of a round bar test piece having a diameter of 16 mm, and frictional sliding of a JIS S45C pin heated to about 800 ° C. was repeated to generate a heat crack on the end face of the test piece. Causes accompanying wear. The heating temperature of the test piece was 600 ° C, and the sliding with the pin was at a surface pressure of 14 N / mm. ² Was performed at a rotation speed of 400 rpm for 5 seconds, and cooling was performed with water at 32 ° C. for 3.5 seconds. After repeating 2,000 cycles of heating, friction sliding and cooling as one cycle, the wear depth of the test piece end face was measured with a surface roughness meter to determine the amount of wear, and the cross-sectional microstructure of the wear portion of the test piece was observed. The number and maximum length of heat cracks generated by the heat treatment were measured.
[0036]
Furthermore, in order to evaluate the oxidation resistance, a round bar test piece having a diameter of 8 mm and a length of 15 mm was prepared from the alloys shown in Table 1, and held in the atmosphere at 900 ° C. for 16 hours, followed by air cooling to room temperature. The change in oxidation weight after 5 repetitions was measured.
Table 2 shows the results of the various tests described above. It was confirmed that the alloy of the present invention had a hardness of 700 ° C. and a tensile strength of 700 ° C. higher than those of the comparative alloy and the conventional alloy, and had excellent mechanical properties in a high temperature range. The thermal expansion coefficient of the conventional alloy No. A value close to that of hot tool steel represented by No. 54 was shown. Conventional alloy No. 52 is a conventional alloy No. 52. It can be seen that the coefficient of thermal expansion is higher than that of hot tool steel represented by No. 54.
The change in oxidized weight after repeated 5 times of holding at 900 ° C. for 16 hours in the atmosphere is as follows: Except for 53 and 54, all others showed increased values, and the adhesion of the oxide film was good.
[0037]
[Table 2]

[0038]
However, the comparative alloy No. When the amount of Cr is low, as in the case of Alloy No. 42, the amount of increase is large. When the Cr content is further decreased as in the case of 51, the increase value is obviously larger than that of the alloy of the present invention, and it is understood that the improvement in oxidation resistance requires the same Cr content as the alloy of the present invention. On the other hand, the conventional alloy No. Hot tool steels such as 53 and 54 exhibited weight loss values. This is because the oxide film was peeled off during the oxidation test, and the oxidation weight was calculated excluding the peeled oxide film. It can be seen that the adhesion of the oxide film was inferior to that of the alloy of the present invention.
[0039]
In addition, FIG. In order to explain the microstructure and the microstructure of the surface portion of the cross section after the oxidation test of No. 10 by a scanning electron microscope, schematic diagrams of constituent phases based on observation are shown. The phases constituting the oxide film were identified by microscopic X-ray diffraction and energy dispersive X-ray analysis. In the high-temperature oxidizing atmosphere of 700 ° C. to 1000 ° C., the alloy of the present invention has a surface oxide layer mainly composed of Cr, a granular α-W layer, and an internal oxide layer mainly composed of Al and Ti. It has a three-layer structure. That is, the granular α-W layer is sandwiched between the surface oxide layer mainly composed of Cr and the internal oxide layer mainly composed of Al and Ti. This is the reason why they have high oxidation resistance and oxide film adhesion, and high hot slidability and hot wear resistance associated with them.
[0040]
In the hot sliding test, the alloy of the present invention was the same as the conventional alloy No. It was confirmed that the seizure load was higher than that of hot tool steel represented by No. 54. Hot slidability requires the formation of an oxide film having excellent adhesion and stability and high-temperature proof stress. The alloy of the present invention satisfies these requirements and has been confirmed to be excellent in hot slidability.
In the hot abrasion test, the alloy of the present invention showed a wear amount of the conventional alloy No. The result was much less than that of hot tool steel represented by No. 54. Conventional alloy No. W containing a large amount of W No. 51 has a high seizure load in the hot sliding test, but the α-W content exceeds 30% by volume and the Cr content is low. Due to lack of stability, it is inferior in wear resistance as compared with the alloy of the present invention. Conventional alloy No. Except for 53, a heat crack occurred. Conventional alloy No. 53 is AC ₁ Due to the low point, the temperature of the end face of the test piece easily exceeds this temperature due to frictional sliding and becomes austenite, and the strength of the frictional sliding part is extremely reduced. It is not observed later.
[0041]
Conventional alloy No. No. 52 had a high thermal expansion coefficient, so that a heat crack was easily generated. Therefore, the heat crack increased the coefficient of friction of the friction sliding surface, and the wear was larger than that of the alloy of the present invention. Comparative alloy No. with high C content Reference numeral 41 indicates a chain structure of MC type carbide, and the portion is a starting point of a heat crack. Comparative alloy no. In No. 42, the tensile strength at 700 ° C. is lower than that of the alloy of the present invention, and the abrasion is large because the γ ′ content is lower than 10% in atomic%. This test confirmed that the alloy of the present invention was excellent in hot wear resistance with heat cracks.
[0042]
(Example 2)
The alloy No. 1 of the present invention shown in Table 1 was used. 4 and Conventional Alloy No. 4 The same heat treatment as in Example 1 was performed for 52 and 53, and a machinability test using a ball end mill was performed. Using a machining center, a ceramic coated carbide ball end mill was used, and a water-soluble cutting oil was used. The test conditions were a cutting speed of 50 m / min, a feed of 0.05 mm / blade, a cutting depth of 2 mm in an axial direction, 1 mm in an outer peripheral blade direction, and a cutting length of 25 m, and a flank wear width of a cutting edge of an end mill was measured. . Table 3 shows the test results.
The alloy of the present invention has the machinability of the conventional Fe-based super heat-resistant alloy No. Needless to say, the conventional hot tool steel No. Excellent machinability was also exhibited as compared with 53. This good machinability is provided by the dispersion of the fine α-W phase. As described above, since the alloy of the present invention has sufficient machinability even after aging treatment, it can be used as a pre-hardened tool that does not require heat treatment after die working.
[0043]
[Table 3]

[0044]
(Example 3)
Hereinafter, an example in which the alloy of the present invention is applied to a hot forging die will be described.
The alloy No. 1 of the present invention shown in Table 1 was used. 10, no. No. 14 and a similar alloy No. 14 and the conventional alloy No. 14 A material having a composition of 54 was prepared, a hot forging die was manufactured from this material, and a practical test was conducted. The results are shown in Table 4. The mold is a mold for manufacturing a gear shaft of an automobile part, and has a size of 80 mm in diameter × 160 mm in length. The mold for manufacturing the gear shaft was composed of a rough land mold for forming a rough ground and a finish mold for finish molding, and a test was performed using a rough ground mold having severe wear. The work material is S35C, which is heated to 1200 ° C. by a high frequency heating device. Forging was performed using a crank press having a maximum capacity of 1600 t, and a white lubricant was sprayed on the surface of the mold after each stamping.
Conventional alloy No. 54 is JIS SKD61. In the heat treatment, after roughing into a hot forging die, the material was heated to 1020 ° C., oil quenched by immersion in 200 ° C. oil, and tempered to a hardness of 446 HV. After that, finish processing was performed, and it was subjected to a practical test.
[0045]
[Table 4]

[0046]
Table 4 shows the life of hot forging dies made of these alloys. Due to this hot forging work, the mold surface is greatly affected by the contact of the mold with the high-temperature work and the sliding of the work. With the change to the lubricant, the effect of the sliding heat generation with the work in particular has increased. No. 54 is a softening of martensite and AC ₁ Austenite transformation was caused by exceeding the transformation point, the strength of the surface layer was reduced, abrasion was likely to occur, and the life was shortened early. Further, since the white lubricant is sprayed on the mold surface after each stamping, a heat cycle of heating and cooling is applied to the mold surface, and heat cracks are generated on the mold surface. Heat cracks generated on the mold surface increase the coefficient of friction with the work, increasing the amount of heat generated by sliding on the work, increasing the thermal effect on the mold surface, and increasing the shear stress generated on the mold surface. Since the plastic flow of the surface layer of the mold increases, abrasion is promoted. Conventional alloy No. In No. 54, the heat effect was increased due to the occurrence of heat cracks, and the generated shear stress exceeded the strength of the surface layer of the mold, and abrasion was likely to occur, and the life was shortened early.
[0047]
The alloy No. of the present invention. 10, No. No. 14 suppresses the heat generated by sliding with the work due to its excellent oxide film properties, and the high-temperature strength surpasses the shear stress generated by the forging operation even when affected by heat, so that the plastic flow of the surface layer No abrasion occurred, wear did not easily occur, and the life was extended.
The alloy No. of the present invention. 10, no. 14 also had a heat crack. No heat crack. 14 are more frequently generated and have a shorter life. No. 14 is an alloy having a large amount of α-W. As a result of observing the microstructure of the wear portion of the mold after the practical test, heat cracks are generated along the α-W, and the amount of α-W is large. In this point, heat cracks were easily generated and abrasion increased. In this practical test, the occurrence of heat cracks is closely related to wear, but when applied to a mold with insufficient lubrication, the effect of anti-seizure during hot sliding of α-W was exhibited.
[0048]
The alloy of the present invention is superior to the conventional alloy in the hardness when heated, but may be inferior to the conventional alloy in the hardness at room temperature. In such a case, in a thick mold, the mold may be deformed by the load of the press during the forging operation. However, the alloy of the present invention is used for the engraved portion, and the conventional alloy No. is used for the base portion. The effect of improving the wear resistance of the alloy of the present invention was exhibited by using a mold combining S.54 and JIS SKT4.
From the above, it has been clarified that the alloy of the present invention has excellent performance as a hot forging die which is often operated mainly in a temperature range of about 700 to 900 ° C.
[0049]
(Example 4)
Hereinafter, an example in which the alloy of the present invention is applied to a Cu alloy extrusion tool will be described.
The alloy No. 1 of the present invention shown in Table 1 was used. No. 10 which had been re-melted with similar components, and Conventional Alloy No. A material having a composition of 52 was prepared, a hot-extrusion tool was manufactured from this, and a practical test was performed. The results are shown in Table 5. Conventional alloy No. shown in Table 1. 52 is JIS SUH660, which is known as A286. This alloy is known as a hot extrusion tool for Cu or Cu alloy.
[0050]
[Table 5]

[0051]
A hot extrusion tool having a double structure by shrink fitting is used, JIS SKT4 is used for the outer cylinder, and the inner cylinder is made of the alloy No. of the present invention. 10 and conventional alloy no. 52 were manufactured and compared. The outer cylinder has an outer diameter of 200 mm, the inner cylinder has an outer diameter of 100 mm, and an inner diameter of 60 mm. 10 and conventional alloy no. 52 were manufactured. An extrusion test of a pure copper billet at 950 ° C. was performed by a 100 t press using these tools. The inner cylinder has a high temperature of about 800 ° C and 500 N / mm ² Exposure to high and low pressures causes a turtle-shaped heat crack due to thermal stress, causing surface peeling, resulting in a life.
Conventional alloy No. In the case of No. 52, the occurrence of heat cracks was already observed on the inner diameter surface after about 10,000 moldings. In the case of 10, a slight heat crack was recognized after about 30,000 moldings.
The alloy No. of the present invention. No. 10 is a conventional alloy No. Compared with No. 52, the tensile strength in a high temperature range is large, and the thermal expansion coefficient is small, so that heat cracks are hardly generated, which drastically improved the life of the extrusion tool. From this result, it became clear that the alloy of the present invention has excellent performance as a hot extrusion tool.
[0052]
【The invention's effect】
The alloy of the present invention is an alloy excellent in hot sliding properties, hot wear resistance, high temperature properties such as heat crack resistance and oxidation resistance, and machinability and hot workability, and is used for automobiles. When applied to hot and hot forging dies for molding automotive parts such as crankshafts and connecting rods, and hot extruded tools of Cu, Al or their alloys, the life of the tool is dramatically improved compared to conventional materials. Can be improved.
[Brief description of the drawings]
FIG. 1 shows the alloy No. 1 of the present invention. 10 is a diagram showing a metal microstructure photograph of a surface portion of a cross section by a scanning electron microscope and a schematic diagram of constituent phases after air-cooling treatment is repeated 5 times to room temperature after holding in air at 900 ° C. for 16 hours. .

Claims

It has a structure that has been subjected to aging treatment, and the phase constituting the structure is 0.02-1.5% in terms of the amount of MC type carbide calculated from the atomic% by the formula (1), and calculated from the atomic% by the formula (2). and 0.5 to 30% by alpha-W volume gamma prime (gamma ') amount 10 to 30%, and represented by the volume% being, Ri Do austenite phase and the balance mainly Ni, alloys The composition is, by weight%, C 0.002 to 0.15%, Si 2% or less, Mn 3% or less, Cr more than 10% and 25% or less, W 10 to 30%, Fe 15% or less, Al 0.4 to 2.5%, comprises Ti0.4~3.5%, the balance being impurities and Ni Tona Rukoto unavoidable Ni-base superalloy.
(1) 2 [C]
(2) 4 (0.026 [Cr] +0.13 [Mo] +0.13 [W] +0.61 [Al] +0.68 [Ti] +0.5 [Nb] +0.5 [Ta]-[C ])
(However, in the above formula, the element without addition is calculated as 0)

The Ni-base superalloy according to claim 1 , wherein a part of Ni is replaced by one or two kinds of Nb of 3.0% or less and Ta of 3.0% or less by weight.

In weight percent, of 10% or less of Mo, W + 2Mo ≦ 30 Ni-base superalloy according to claim 1 or 2 was replaced with a portion of the Ni in the range of.

The Ni-based super heat-resistant alloy according to any one of claims 1 to 3, wherein a part of Ni is replaced by one or two kinds of Zr of 0.2% or less and B of 0.02% or less by weight%.

The Ni-base superalloy according to any one of claims 1 to 4, wherein a part of Ni is replaced by one or two of 0.02% or less of Mg and 0.02% or less of Ca by weight.