JP3559670B2

JP3559670B2 - High-strength Ni-base superalloy for directional solidification

Info

Publication number: JP3559670B2
Application number: JP02473697A
Authority: JP
Inventors: 英樹玉置; 明吉成; 昭岡山; 満小林; 景弘影山; 丈博大野
Original assignee: Hitachi Ltd; Hitachi Metals Ltd
Current assignee: Hitachi Ltd; Hitachi Metals Ltd
Priority date: 1996-02-09
Filing date: 1997-02-07
Publication date: 2004-09-02
Anticipated expiration: 2017-02-07
Also published as: JPH09272933A

Description

【０００１】
【発明の属する技術分野】
本発明は、ガスタービン等の高温機器の動翼，静翼等の高温部品に用いられる新規なＮｉ基超合金に関し、特に、優れた高温強度を有し、高い歩留まりで生産することが困難な、大型複雑形状を有する単結晶合金からなる高温部材に最適なＮｉ基超合金に関する。
【０００２】
【従来の技術】
ガスタービンの燃焼ガス温度は、熱効率の向上を目的に年々上昇する傾向にあり、ガスタービンの各高温部材には従来より高温強度の優れた材料が必要とされている。そのため、ガスタービンの高温部材中で最も苛酷な環境に曝される動翼用の材料は、Ｎｉ基超合金の普通鋳造材から柱状晶材へと変遷し、さらに航空機エンジン用のガスタービンでは、より高温強度の優れた単結晶材が実用化されている。ここで、柱状晶材及び単結晶材は方向性凝固材の一種であり、どちらも一方向凝固法と呼ばれる方法で鋳造される。柱状晶材では、米国特許第３，２６０，５０５号公報等に示される方法で結晶粒を一方向に細長く成長させ、主応力の作用方向に垂直な粒界を極力少なくすることで高温強度の向上が可能となる。また、単結晶材では、米国特許第３，４９４，７０９号公報等に示される方法により、鋳物全体を一つの結晶とすることでより一層の高温強度の向上が可能となる。
【０００３】
さらに、Ｎｉ基超合金の高温強度を向上させるためには、析出強化相である γ′相を均一微細に析出させる溶体化熱処理が有効である。すなわち、これらのＮｉ基超合金はＮｉ_３（Ａｌ，Ｔｉ，Ｎｂ，Ｔａ）を主体とするγ′相の析出により強化されるが、このγ′相は均一微細に析出していることが好ましい。ところが、凝固のままの状態では粗大なγ′相（凝固後の冷却中に析出し粗大化した γ′相と最終凝固部に粗大に晶出した共晶γ′相）が存在するため、一度高温に加熱してγ′相を基地のγ相に固溶させた後急冷し（溶体化熱処理）、それに続く時効処理中に均一微細なγ′相として析出させることにより高温強度を向上することが可能となる。この溶体化熱処理は、γ′相の固溶温度以上で、かつ初期溶融温度以下のできるだけ高い温度で行うことが好ましい。これは、熱処理温度が高くなるほどγ′相が均一微細化される領域が多くなり、さらに、γ′相が均一微細化された領域が増加するほど高温強度が向上するためである。単結晶材の高温強度が優れているもう一つの理由は、初期溶融温度を著しく低下させる結晶粒界強化元素を不純物程度にしか含まない単結晶専用の合金を用いることで、溶体化熱処理温度を高くすることが可能となり、その結果、凝固後に粗大に析出したγ′相のほとんど全てを均一微細化できるためである。
【０００４】
上記のとおり、ガスタービンの動翼材として現状の技術で最も高温強度に優れているのはＮｉ基超合金の単結晶材であり、そのための合金としてＣＭＳＸ−４（米国特許第４，６４３，７８２号公報），ＰＷＡ１４８４（米国特許第４，７１９，０８０号公報）及びＲｅｎｅ’Ｎ５（特開平５−５９４７４号公報）等の単結晶合金が開発され、航空機エンジン用のガスタービンの動翼に適用されている。しかし、前述の通り、これらの単結晶合金はいずれもＣ，Ｂ及びＨｆ等の結晶粒界強化元素を不純物レベルにしか含んでいない。従って、単結晶合金を用いて鋳造した動翼中に結晶粒界が存在すると、強度が極端に低下し、場合によっては凝固時に既に結晶粒界に沿った縦割れが生じてしまう。そのため、単結晶合金で鋳造した動翼をガスタービンで使用するためには、動翼全体を完全な単結晶にする必要がある。航空機エンジン用のガスタービンの動翼は、全長が最大でも１００ｍｍ程度であるため、鋳造時に結晶粒界が発生する確率は小さく、単結晶合金でもある程度の歩留まりでの生産が可能である。しかし、発電用ガスタービンの動翼では、全長が約１５０〜４５０ｍｍとなり、動翼全体を完全な単結晶とすることは非常に難しい。従って、現状の技術では、単結晶合金を用いて発電用ガスタービンの動翼を高い歩留まりで生産することは困難である。
【０００５】
上記のように、鋳造歩留まりの関係で単結晶合金が適用できない大型動翼の高温強度向上を目的に、高温強度の優れた柱状晶材用の合金の開発が進められ、その結果、ＣＭ１８６ＬＣ（米国特許第５，０６９，８７３号公報）やＲｅｎｅ’１４２（米国特許第５，１７３，２５５号公報）等の柱状晶合金が開発された。これらの合金は、鋳造時の凝固割れを防止し、使用中の信頼性を確保するのに十分な量の結晶粒界強化元素を含有し、かつ、ＰＷＡ１４８０（米国特許第４，２０９，３４８号公報），ＣＭＳＸ−２（米国特許第４，５８２，５４８号公報）あるいはＲｅｎｅ’Ｎ４（米国特許第５，３９９，３１３号公報）等の第一世代の単結晶合金に匹敵する高温強度を有している。従って、これらの柱状晶合金を用いることで、第一世代の単結晶合金からなる動翼とほぼ同等の高温強度を有する動翼を、より高い歩留まりで生産することが可能となった。しかし、ガスタービンの燃焼ガス温度はさらに向上する傾向にあり、これら従来の柱状晶合金ではガスタービンのさらなる熱効率向上の要求を満たすためには高温強度が不十分となってきた。
【０００６】
また、特開平７−１４５７０３号公報及び特開平５−５９４７３号にはＣ，Ｂ，Ｚｒ及びＨｆを含む柱状晶を含む単結晶合金が開示されている。
【０００７】
以上のことから、発電用ガスタービンの高効率化のためには、従来技術ではあい矛盾する高い生産歩留まりと高温強度を両立した合金の開発が必須である。
【０００８】
【発明が解決しようとする課題】
前述のとおり、Ｎｉ基超合金の高温強度を向上させるためには溶体化熱処理温度をできるだけ高くする方法が有効であり、そのためには結晶粒界強化元素の添加量を不純物程度とすることが好ましい。しかし、一方で、高い生産歩留まりと使用中の信頼性を確保するためには、結晶粒界に適当な強度を付与するための結晶粒界強化元素が合金中に含まれている必要がある。従って、従来は高温強度を向上させるためには結晶粒界の強度を犠牲にする必要があり、反対に、結晶粒界の強度を向上させるためには高温強度を犠牲にする必要があった。
【０００９】
さらに、本発明者らが従来柱状晶合金ＣＭ１８６ＬＣに対して行った研究によれば（材料とプロセス，Ｖｏｌ．７（１９９４），ｐ１７９７及び材料とプロセス，Ｖｏｌ．８（１９９５），ｐ１４５８）、結晶粒界強化元素であるＢは、溶体化熱処理中に結晶粒界から粒内へ拡散してしまうことが明らかにされている。従って、この合金は、結晶粒界強化元素を含んでいるものの、高温強度を向上させるために溶体化熱処理を行うと結晶粒界の強度が実用に耐えられないレベルまで低下してしまう。通常、方向性凝固材においては、主応力の作用する方向が凝固方向に合わされるため、高温強度は凝固方向の強度で評価される。この場合、高温強度、つまり結晶粒界に平行な凝固方向の強度はγ′相の溶体化が進むに従って向上するが、反対に、結晶粒界の強度、つまり結晶粒界に垂直な凝固方向に垂直方向の強度は低下することになる。
【００１０】
以上のことから、既存の単結晶合金に単純に結晶粒界強化元素を添加しても、歩留まりの向上は望めるが、溶体化熱処理温度が著しく低下するため、優れた高温強度は望めない。また、従来の柱状晶合金については、初期溶融及び結晶粒界の強度低下の問題から、溶体化熱処理温度をこれ以上高くすることはできず、現状以上の高温強度の向上は望めない。
【００１１】
本発明の目的は、鋳造時の凝固割れを防止し、さらに使用中の信頼性を確保するのに十分な結晶粒界の強度を有し、かつ従来の柱状晶合金より優れた高温強度をあわせ持つ方向性凝固用高強度Ｎｉ基超合金を提供することにある。
【００１２】
【課題を解決するための手段】
本発明は、従来技術では相反する関係にあると考えられていた、高温強度と結晶粒界の強度を両立する合金組成を得ることを目的に、Ｃ，Ｂ，Ｈｆ及びＺｒの４種の結晶粒界強化元素を様々な組み合わせで単結晶合金に添加し、結晶粒界強化元素の添加量と高温強度，結晶粒界の強度及び溶体化熱処理の関係を調べた結果なされたものである。
【００１３】
単結晶合金の成分に調製した多結晶マスターインゴットに、結晶粒界強化元素を一方向凝固炉内で添加し、目的の組成の一方向凝固平板を鋳造した。ここで、各組成の高温強度は凝固方向のクリープ破断強度で評価した。また、鋳造性と使用中の信頼性を確保するための結晶粒界の強度は平板の凝固方向に垂直な方向、つまり結晶粒界が応力負荷方向に対して垂直になる方向のクリープ破断強度と高温引張特性で評価した。
【００１４】
この結果、従来知られていたＢの最適添加量よりかなり高い領域に、凝固方向と凝固方向に垂直な方向の強度、つまり、高温強度と結晶粒界の強度を両立させる最適添加量が存在することが明らかになった。つまり、結晶粒界強化元素としてＣを０．０３〜０．２０％、好ましくは０．０５〜０．１％未満添加し、Ｈｆが１．５％以下、好ましくは０．０１〜０．５％未満，Ｚｒが０．０２％以下、好ましくは０．０１％未満含まれる場合、凝固方向の強度と凝固方向に垂直な方向の強度の両方に最適なＢ量は０．００４〜０．０５％、好ましくは０．０１５％を超え０．０４％以下の範囲であり、特に０．０３％付近で凝固方向の強度と凝固方向に垂直な方向の強度の両方が最大となることが明らかになった。好ましくは、従来の柱状晶合金のＢの添加量がいずれも約０．０１５％程度であることを考えると、従来合金の約２倍の量である。
【００１５】
Ｂは合金の初期溶融温度を著しく低下させる元素である。従って、Ｂを多量に添加する場合には、初期溶融温度の低下の影響も考慮する必要があるが、本発明合金においては従来合金の約２倍の量のＢを添加した組成においても著しい初期溶融温度の低下は見られなかった。
【００１６】
また、Ｃも高温強度と結晶粒界の強度の両立を図るうえで重要な元素である。つまり、結晶粒界強化元素としてＢを０．００７〜０．０１５％，Ｈｆを０．５％未満，Ｚｒを０．０１％未満含む合金において、凝固方向のクリープ破断強度はＣの添加量が増えるに従って低下することが明らかになった。これに対して、凝固方向に垂直方向のクリープ破断強度は、０．２０％以下，好ましくは０．１０％までは添加量が増えるに従って向上し、それ以上の添加量では、０．１０％をピークにして低下することがわかった。従って、凝固方向に垂直方向のクリープ破断強度のみを考慮すれば、Ｃの最適な添加量は約０．１％付近に存在するが、一方で凝固方向のクリープ破断強度はＣ量の増加にともなって低下することを考慮すれば、高温強度と結晶粒界の強度を両立するために最適なＣ量の範囲は０．０５〜０．１％未満の範囲であると考えられる。Ｃの添加量が０．０３％未満、より０．０５％未満である場合は、高温強度は優れているが結晶粒界の強度が低く、鋳造時の凝固割れを防止し、使用中の信頼性を確保することができない。一方、Ｃの添加量が０．２％以上、好ましくは０．１％以上になると、高温強度が著しく低下し、かつ結晶粒界の強度も低下してしまう。
【００１７】
Ｚｒ及びＨｆは同族の元素であり、Ｎｉ基超合金に及ぼす影響もほぼ同等であることが知られている。本発明における検討では、Ｚｒは合金の初期溶融温度を著しく低下させ、高温での溶体化熱処理を不可能とすることで、合金の凝固方向のクリープ破断強度を低下させることが明らかになった。さらに、Ｚｒは横方向のクリープ破断強度についても効果が無いことが明らかになったことから、Ｚｒの添加量は０．０１％未満とすることが望ましく、より好適にはＺｒを実質的に無添加とする必要がある。ＨｆはＺｒと同様に、合金の初期溶融温度を著しく低下させ、高温での溶体化熱処理を不可能とすることで、合金の凝固方向のクリープ破断強度を低下させ、かつ横方向のクリープ破断強度についてもほとんど効果が見られない。しかし、Ｈｆは横方向の引張延性を改善する効果があり、さらに、０．２５％程度の添加量では、凝固方向のクリープ破断強度を若干低下させるものの、横方向のクリープ破断強度と引張り強さの両方を改善することが明らかになった。従って、Ｈｆの添加量は０．０１〜０．５％未満とすることが望ましく、好適な添加量の範囲は０．２〜０．４％である。さらに、高温強度と結晶粒界の強度を両立するための最適な添加量の範囲は０．２〜０．３％である。
【００１８】
本発明合金においては、以上のとおりＢ，Ｃ，Ｈｆ及びＺｒの添加量の組み合わせを最適化することで、結晶粒界の強度を得るために十分な結晶粒界強化元素を含有し、かつ高温強度を向上させるために十分な溶体化熱処理を行うことを可能とすることで、従来技術では成しえなかった高温強度と結晶粒界の強度を両立することが可能となった。
【００１９】
また、上記の結果は結晶粒界強化元素の組み合わせのみから決まるわけではなく、結晶粒内の強化に寄与する元素の効果も無視できない。本発明合金の組成における結晶粒界強化元素以外の特徴の一つは、Ｃｏの添加量にある。従来の柱状晶合金には９％以上の多量のＣｏを含有している合金が多いが、本発明においては、Ｃｏの多量な添加は合金の高温強度を著しく低下させ、かつ結晶粒界の強度にはほとんど影響を及ぼさないことが明らかになった。しかし、一方で、Ｃｏは燃焼ガス雰囲気中での耐食性を改善する効果があり、耐食性が重視される発電用ガスタービン用動翼及び静翼の材料には、高温強度を著しく低下させない範囲での添加が必須である。
【００２０】
さらに、本発明合金において、溶体化熱処理により結晶粒界の強度を低下させることなく高温強度を向上させることができたのは、Ｔａの添加量を最適化できたことも大きな要因である。上述のとおり、従来柱状晶合金のＣＭ１８６ＬＣにおいては、高温強度を向上させるために溶体化熱処理を行うと結晶粒界からＢが結晶粒内へ拡散してしまい、結晶粒界の強度が極端に低下する。これは、溶体化熱処理でγ′相をγ相中に一旦固溶させる際に、結晶粒界近傍のγ′相がγ相中に固溶するのとほぼ同時にＢがγ相中に拡散を開始し、最終的には結晶粒界からＢが消失してしまうためである。この課題を解決するために、本発明合金では従来合金と比べて著しく高い量のＴａを添加した。この結果、結晶粒界近傍のγ′相の固溶温度が結晶粒内と比べてかなり高くなり、結晶粒界近傍のγ′相をγ相中に固溶させることなく、結晶粒内のγ′相を溶体化することが可能となった。従って、本発明合金においては、結晶粒界からＢを拡散消失させることなく結晶粒内の強度を向上させることができ、その結果、結晶粒界の強度を低下させることなく高温強度を向上させることが可能となった。
【００２１】
一般に、溶体化率、つまり合金中でγ′相が微細化されている領域の割合が増えるほど高温強度が向上するが、上記のとおり結晶粒界の強度を考慮すれば、溶体化率はできるだけ少ないことが好ましい。従って、合金の高温強度と結晶粒界の強度を両立させるためには、溶体化率が少ない場合でも優れた高温強度を示すことができる組成が好ましい。そこで、本発明合金では固溶強化に有効なＲｅ及びＷの添加量を最適化し、合金を最大限固溶強化することで、比較的低い溶体化率で高温強度を向上させることを可能としている。
【００２２】
本発明の合金は、一方向凝固法により方向性凝固して使用するのに好適である。特にガスタービン用の動翼においては、遠心力が作用する方向を凝固方向として鋳造することが望ましい。また、これまではガスタービン用の動翼での使用を前提に述べてきたが、ガスタービン用の静翼等の他の高温部品に用いることもできる。ガスタービン用の静翼の場合には、熱応力が最大となる方向に凝固方向をあわせて使用することが好ましい。本発明合金は、通常の柱状晶動翼，柱状晶静翼に使用できるのはもちろんのこと、単結晶鋳造中に動翼の一部分に結晶粒界が発生した動翼に使用することも可能である。このような動翼は従来は不良品とされていたが、本発明合金を用いればこのような動翼でも十分使用可能であり、その結果、単結晶動翼の鋳造歩留まりを大幅に改善することが可能となる。また、本発明合金は通常の単結晶動翼に用いることもできる。従来単結晶合金を用いて高い歩留まりで完全な単結晶動翼，単結晶静翼を鋳造することができる場合でも、本発明合金では結晶粒界の有無の検査を簡素化することが可能なため、生産コストを低減することが可能である。さらに、従来は抜き取り破壊試験で動翼内面の結晶粒界の有無を保証していたが、本発明合金は結晶粒界が存在しても強度が保証できるため、動翼の信頼性を大幅に向上することが可能となる。
【００２３】
以上本発明は、重量で、Ｃ：０．０３〜０．２０％、好ましくは０．０５％以上０．１％未満，Ｂ：０．００４〜０．０５％、好ましくは０．０１５％を越え０．０４％以下、Ｈｆ：１．５％以下、好ましくは０．０１〜０．５％未満，Ｚｒ：０．０２％以下、好ましくは０．０１％未満，Ｃｒ：１．５〜１６％，Ｍｏ：６％以下，Ｗ：２〜１２％，Ｒｅ：０．１〜９％，Ｔａ：２〜１２％，Ｎｂ：４．０以下、好ましくは０．３〜４％，Ａｌ：４．０〜６．５％，Ｔｉ：０．４％未満、好ましくは添加せず、Ｃｏ：９％以下及び６０％以上のＮｉを含む結晶粒界強度の優れた方向性凝固用高強度Ｎｉ基超合金である。特に高温強度と結晶粒界の強度を高度に両立し、かつ燃焼ガス中で優れた耐食性を示す合金は、重量％で、Ｃ：０．０５〜０．１０％，Ｂ：０．０１８〜０．０４％，Ｈｆ：０．０１〜０．５％未満，Ｚｒ：０．０１％未満，Ｃｒ：４〜１２．５％，Ｍｏ：４．５％以下，Ｗ：５〜１０％，Ｒｅ：１〜６％，Ｔａ：５〜１２％，Ｎｂ：０．３〜３％，Ａｌ：４．０〜６．０％，Ｃｏ：０．５〜１．２％及び残部不可避の不純物とＮｉからなる結晶粒界強度の優れた方向性凝固用高強度Ｎｉ基超合金である。さらに、より高温強度の優れた合金が要求される場合には、重量％で、Ｃ：０．０６〜０．１０％，Ｂ：０．０１８〜０．０３５％，Ｈｆ：０．１〜０．５％，Ｃｒ：６．５〜８．５％，Ｍｏ：０．４〜３．０％，Ｗ：５．５〜９．５％，Ｒｅ：１．０〜６．０％，Ｔａ：６〜１０．５％，Ｎｂ：０．３〜１．５５％，Ａｌ：４．０〜６．０％，Ｃｏ：０．５〜２．５％及び残部不可避の不純物とＮｉからなる結晶粒界強度の優れた方向性凝固用高強度Ｎｉ基超合金が適当であり、さらに好適な組成は、重量％で、Ｃ：０．０６〜０．１０％，Ｂ：０．０１８〜０．０３５％，Ｈｆ：０．２〜０．３％，Ｃｒ：６．９〜７．３％，Ｍｏ：０．７〜２．０％，Ｗ：７．０〜９．０％，Ｒｅ：１．２〜２．０％，Ｔａ：８．５〜９．５％，Ｎｂ：０．６〜１％，Ａｌ：４．０〜６．０％，Ｃｏ：０．５〜１．２％，Ｔｉ０．４％未満、好ましくは添加せず及び６０％以上のＮｉ、好ましくは残部不可避の不純物とＮｉからなる組成である。
【００２４】
また、特に燃料中にＳ等の不純物が多い環境においては、重量％で、Ｃ：０．０６〜０．０８％，Ｂ：０．０１８〜０．０３５％，Ｈｆ：０．２〜０．３％，Ｃｒ：６．９〜７．３％，Ｍｏ：０．７〜１％，Ｗ：８〜９％，Ｒｅ：１．２〜１．６％，Ｔａ：８．５〜９．５％，Ｎｂ：０．３〜１％，Ｔｉ：０．５％未満，Ａｌ：４．９〜５．２％，Ｃｏ：０．８〜１．２％及び残部不可避の不純物とＮｉからなることを特徴とする結晶粒界強度の優れた方向性凝固用高強度Ｎｉ基超合金が適当である。
【００２５】
以上の組成からなるＮｉ基超合金を用いれば、凝固方向のクリープ破断寿命が１０４０℃，１４ｋｇｆ／ｍｍ^２の条件で３５０時間以上、かつ凝固方向に垂直な方向のクリープ破断寿命が９２７℃，３２ｋｇｆ／ｍｍ^２の条件で３０時間以上である、高温強度と結晶粒界の強度の両方に優れた方向性凝固鋳物を得ることができる。
また、上記の組成からなるＮｉ基超合金を用いれば、溶体化熱処理により体積率で５０％以上の領域でγ′相を一辺が０．５μｍ以下の直方体状に整えることが可能で、かつ凝固方向に垂直な方向のクリープ破断寿命が９２７℃，３２ｋｇｆ／ｍｍ^２の条件で３０時間以上、凝固方向に垂直な方向の８００℃での引張り強さが９５ｋｇｆ／ｍｍ^２以上の特性を有する、高温強度と結晶粒界の強度の両方に優れた方向性凝固鋳物を得ることができる。
【００２６】
本発明は、重量で、Ｃ０．０３〜０．２０％，Ｂ０．００４〜０．０５％，Ｃｒ４．０〜１２．５％，Ｍｏ４．５％以下，Ｗ５．０〜１０．０％，Ｒｅ１．０〜７．０％，Ｔａ５．０〜１２．０％，Ｎｂ０．３〜４．０％，Ａｌ４．０〜６．５％，Ｔｉ０．４％未満，Ｃｏ０．５〜５．０％，Ｈｆ１．５％以下，Ｚｒ０．１５％以下及び６０％以上のＮｉを含み、前記Ｃ量は０．１５より前記Ｂ量に５．４５倍した値を差し引いた値以上であることを特徴とする方向性凝固用Ｎｉ基超合金にある。
【００２７】
特に、Ｃ量とＢ量とは（０．２０％，０．０３％）と（０．０８％，０．０５％）とを結ぶ直線以下、より好ましくは（０．２０％，０．０１％）と（０％，０．０４７％）とを結ぶ直線以下である。
【００２８】
本発明は、凝固方向の１０４０℃，１４ｋｇ／ｍｍ^２でのクリープ破断時間が３５０時間以上及び凝固方向に垂直な方向の９２０℃，３２ｋｇ／ｍｍ^２でのクリープ破断時間が３０時間以上であることを特徴とする柱状晶Ｎｉ基超合金にある。特に、前者が５００時間以上、後者が４５時間以上のものが好ましい。
【００２９】
本発明は、凝固方向の１０４０℃，１４ｋｇ／ｍｍ^２でのクリープ破断時間が３５０時間以上及び凝固方向に垂直な方向の９２０℃，３２ｋｇ／ｍｍ^２でのクリープ破断時間が前記凝固方向のクリープ破断時間を０．１５倍した値より３２．５を差し引いた値以上であることを特徴とする柱状晶Ｎｉ基超合金にある。特に、前者が５００時間以上とするものが好ましい。
【００３０】
更に、本発明はＭｏ量に対するＣｏ量の比率が０．２〜５となるようにするのが好ましく、より０．４〜２．０が好ましい。
【００３１】
表１は本発明に係る各合金組成について広い範囲，望ましい範囲，好適範囲，最適範囲について示したものである。
【００３２】
前述した本発明に係るＮｉ基合金は基地のγ相が単結晶からなるものが好ましい。
【００３３】
【表１】

【００３４】
【発明の実施の形態】
〔実施例１〕
表２はＣ，Ｂ，Ｈｆ及びＺｒの結晶粒界強化元素を添加し、その結晶粒界強化元素添加量と高温強度及び結晶粒界の強度の関係を示したものである。ここでベース合金には、重量％で７．８Ｃｒ−７．２Ｗ−１．８Ｍｏ−４．７Ａｌ−１．６Ｎｂ−７．５Ｔａ−１．６Ｒｅ−残部Ｎｉからなるものである。これらの合金は、真空誘導溶解で作製したベース合金の多結晶マスターインゴットに、一方向凝固炉中で結晶粒界強化元素を添加し、１５ｍｍ×１００ｍｍ×１００ｍｍの一方向凝固した柱状晶を有する平板を鋳造した。次にこの一方向凝固平板から１０ｍｍ×１０ｍｍ× １０ｍｍのブロックを数個採取し、１２５０，１２６０，１２７０，１２８０，１２９０，１３００，１３１０，１３２０，１３３０℃の各温度で２時間ずつ熱処理し、熱処理後の組織から各組成に最適な溶体化熱処理条件を決定した。なお、一部の合金については、必要に応じてより細かい範囲で温度を変化させ、より正確な最適溶体化処理条件を検討した。ここで、最適な熱処理条件とは、体積率で５０％以上の領域でγ′相を一辺が０．５μｍ以下の直方体状に整えることができ、かつ初期溶融温度以下のできるだけ高い温度である。この予備検討から決定し、実際に各合金に施した溶体化熱処理の条件を表２中に示す。なお、溶体化熱処理後は空冷とし、これに続く時効熱処理の条件は、全ての合金で１０８０℃／４時間／空冷＋８７１℃／２０時間／空冷の一定とした。
【００３５】
ここで、高温強度は、一方向凝固平板の凝固方向に採取した試料の９２０℃，３２ｋｇｆ／ｍｍ^２の条件のクリープ破断強度で評価した。以後、これを凝固方向のクリープ破断強度と記すこととする。また、結晶粒界の強度は一方向凝固平板の凝固方向に対して垂直方向（以後、これを横方向と記す）、つまり応力軸が結晶粒界に対して垂直方向になるように採取した試料の９２０℃，３２ｋｇｆ／ｍｍ^２の条件のクリープ破断強度及び８００℃での高温引張特性で評価した。これらの結果を表２中に併せて示す。なお、試験片形状は、クリープ破断試験，高温引張試験共に、直径６ｍｍ，標点間距離３０ｍｍとした。これらの試験片は凝固方向と同じ方向では全体が単結晶であるものと同じ特性を示すものと思われる。
【００３６】
なお、一方向凝固平板の結晶粒の幅は、凝固開始側（下部）では約１〜５ｍｍ，上部では約５〜１０ｍｍである。横方向の強度評価用の試料は平板中央部（結晶粒の幅が約５ｍｍ）で採取した。従って、ｇａｕｇｅｌｅｎｇｔｈ内に約５個の結晶粒界が存在することになる。凝固方向の試料採取位置については特に考慮しなかった。極端な場合にはｇａｕｇｅｌｅｎｇｔｈ内が単結晶であった可能性もあるが、通常は３個程度の結晶粒界が存在する。
【００３７】
【表２】

【００３８】
図１はＣを約０．１％含み、Ｈｆ及びＺｒを実質的に無添加とした場合のＢ量と凝固方向及び横方向のクリープ破断強度の関係を示す。この場合、凝固方向及び横方向の両方において、最適添加量が０．０３％付近に存在している。従来の柱状晶合金のＢの添加量が０．０１５％程度のレベルであることを考えると、この結果は、従来最適と考えられてきたＢの添加量の約２倍の量の付近に実際の最適値が存在していたことを示している。Ｂ量は０．０１７〜０．０４０％で高い強度が得られる。
【００３９】
図２はＢを約０．０１％含み、Ｈｆ及びＺｒを実質的に無添加とした場合のＣ量と凝固方向及び横方向のクリープ破断強度の関係を示す。また、図３にＨｆ及びＺｒに加えて、Ｃを実質的に無添加とした場合のＢ量と凝固方向及び横方向のクリープ破断強度の関係を示す。これらの関係から、Ｃの添加により凝固方向のクリープ破断強度は低下するが、Ｃは横方向の強度を得るために必須の添加元素であることがわかる。従って、高温強度と結晶粒界の強度を両立させるためにはＣ量を精密に制御する必要がある。また、Ｃ量を制御することで、高温強度あるいは結晶粒界の強度のどちらかを重視した合金を得ることも可能である。具体的には、高温強度を重視する場合にはＣ量は低いほど良く、高温強度より結晶粒界の強度を重視する場合にはＣ量は多いほうが良い。
【００４０】
図４，図５にＣを約０．１％，Ｂを約０．０１％含む場合のＺｒ及びＨｆ量と凝固方向及び横方向のクリープ破断強度の関係を示す。この結果から、Ｚｒ及びＨｆは凝固方向のクリープ破断強度を低下させ、さらに、横方向のクリープ破断強度をほとんど改善しないことが明らかになった。しかし、Ｈｆは、図６に示すとおり、横方向の引張延性を改善する効果がある。
【００４１】
〔実施例２〕
表３に示す組成の合金の多結晶マスターインゴットを真空誘導溶解で作製し、一方向凝固炉で１５ｍｍ×１００ｍｍ×２２０ｍｍの一方向凝固平板を鋳造した。これらの合金について、Ｎｏ．１〜２５と同様に、一方向凝固平板から１０ｍｍ×１０ｍｍ×１０ｍｍのブロックを数個採取し、１２５０，１２６０，１２７０，１２８０，１２９０，１３００，１３１０，１３２０，１３３０℃の各温度で２時間ずつ熱処理し、溶体化熱処理条件の予備検討を行った。この結果をもとに、これらの合金には、１２５０℃／４時間から順次１０℃／４時間刻みで熱処理温度を上げ、最終的に表２中に記載の最高溶体化熱処理温度で４時間熱処理後空冷とする多段溶体化熱処理を施した。溶体化熱処理に続く時効熱処理は、全ての合金について１０８０℃／４時間／空冷＋８７１℃／２０時間／空冷とした。
【００４２】
表３に各合金の特性評価結果を併せて示す。これらの試験のうち、クリープ破断試験及び高温引張試験用の試料は、Ｎｏ．１〜２５と同じ方法で採取し、試験片形状も同一とした。条件は、凝固方向のクリープ破断試験を温度１０４０℃で応力１４ｋｇｆ／ｍｍ^２，横方向のクリープ破断試験を温度９２７℃で応力３２ｋｇｆ／ｍｍ^２，横方向の引張試験を温度８００℃とした。
【００４３】
また、一部の合金について実機模擬燃焼試験で耐食性を評価した。試験片は直径９ｍｍ，長さ５０ｍｍの丸棒とし、実機模擬燃焼ガス雰囲気中で、９００℃／７ｈ／空冷×７回試験後の脱スケール後の腐食減量を測定した。
【００４４】
【表３】

【００４５】
表に示すように、Ｔａは高温強度を向上するために最低でも２％以上、好ましくは５％以上添加することが望ましく、８．５〜９．５％の範囲に高温強度上の最適値が存在する。一方、Ｔａの大量な添加は前述のとおりγ′相の固溶温度を向上させる。従って、Ｔａを過剰に添加すると、合金の初期溶融温度とγ′相の固溶温度の差が少なくなり、初期溶融を生じることなくγ′相の溶体化できる領域が減少して合金の析出強化量が低下する。そのため、１２％を超えるＴａの添加はもはや高温強度の向上に効果がなく、さらに、上限を１０％以下にすることが好ましい。
【００４６】
Ｃｏは、他の合金元素の添加量がほぼ同一でＣｏの添加量のみが大きく異なるＮｏ．３８及びＮｏ．１００〜１０４の関係から、添加量が増加するに従って高温強度を低下させることが明らかである。従って、高温強度を考慮すれば、Ｃｏの添加量は９％以下、好ましくは９９％未満，より好ましくは０．５〜５％とする。特に、０．５〜１．２％程度のＣｏの添加は耐食性を向上する効果がある。
【００４７】
Ｗ及びＲｅは合金を固溶強化することで高温強度の向上に有効な元素であり、各々２％、好ましくは５％及び０．１％、好ましくは１％以上添加することが望ましい。さらに高温強度を重視する場合には各々５．５％及び１．２％以上添加することが好ましい。一方、これらの元素の効果はある程度の添加量で飽和し、さらに過度に添加すると合金の高温強度はかえって低下してしまう。これは、これらの元素を固溶限を超えて過度に添加すると、主にＷ，Ｒｅからなる針状あるいは板状の析出物が析出するためである。従って、Ｗ及びＲｅの添加量の上限は各々１２％、好ましくは１０％及び９％、好ましくは６％とすることが望ましい。さらに、前記析出物の多量な析出を抑制するためには、Ｗ及びＲｅの添加量を各々９．５％及び３．１％以下とすることが好ましい。本発明合金における、最適なＷの添加量は８．０〜９．０％であり、同じく最適なＲｅの添加量は１．２〜１．６％である。そして、Ｗは５〜１０％，より５．５〜９．５％が好ましく、Ｒｅは１〜６％，より１．２〜１．３％が好ましい。
【００４８】
ＷとＲｅの最適な添加量はＷとＲｅの総量で考えることが好ましく、高温強度はＷ＋Ｒｅ量が９．５〜１２％の範囲で最大となる。これに対して、Ｗ＋Ｒｅ量が９．５％を下回ると固溶強化が不足するため高温強度は低下する。また、Ｗ＋Ｒｅ量が１２％を上回ると、上記の析出物が多量に析出し、特に１０００℃以上でのクリープ強度が大幅に低下する。
【００４９】
ＡｌはＮｉ基超合金の強化因子であるγ′相を形成するための必須元素である。また、表面にＡｌ_２Ｏ_３被膜を形成することで耐酸化性及び耐食性の向上に寄与する。従って、Ａｌの添加量は最低でも４．０％以上、好ましくは４．５％以上であることが望ましい。しかし、６．５％を越えて過度に添加すると合金中の共晶γ′相の量が増加してしまう。本発明合金は、固溶強化に有効な元素の添加量を最適化することで、完全な溶体化処理を行わない状態でも優れた高温強度を示すことができるように考慮されている。従って、共晶γ′相が存在する状態でも優れた高温強度を有する。しかし、クリープ損傷の場合、共晶γ′相は最終的には亀裂の起点となり、材料の破壊時期を早めることになるため、共晶γ′相の量は少ないことが好ましい。従って、Ａｌの添加量は６．５％以下，より５．７％以下とすることが望ましい。特に、４．７〜５．４％，より４．９〜５．２％が好ましい。
【００５０】
ＣｒはＣｒ_２Ｏ_３被膜を形成して合金の耐食性及び耐酸化性を向上する効果があるため、最低でも１．５％以上，より４％以上添加することが望ましい。しかし、Ｃｒを過度に添加すると上述のＷ，Ｒｅからなる析出物の析出を助長するため、高温強度に有効なＷあるいはＲｅの添加量を低減する必要が生じる。従って、高温強度を重視する場合はＣｒの添加量の上限を１６％，より１２．５％とすることが望ましい。特に、６．５〜８．５％，より６．９〜７．３％が好ましい。
【００５１】
ＭｏはＷあるいはＲｅと同様な効果を示すが、合金の燃焼ガス雰囲気中での耐食性を著しく低下させるため、耐食性を重視する場合はＭｏの添加量を６％以下，より４．５％以下に制限することが望ましい。さらに耐食性を重視する場合には、Ｍｏの添加量を０．４〜１％，より０．７〜１％とすることが好ましい。
【００５２】
ＮｂはＴａと同族の元素であり、高温強度についてはほぼ同様の効果があり、０．３〜４％含有される。さらに、硫化物を形成しやすいため、燃料中にＳが多く存在する環境においては、Ｓの合金内部への侵入を遅らせ耐食性を改善する効果がある。しかし、本発明では、合金中に一定量以上のＮｂとＢが存在する場合、共晶部にＮｂとＢと主成分とする低融点相を形成して合金の初期溶融温度を著しく低下させることが明らかになった。この低融点相は、凝固時の偏析により生じるもので、鋳造条件により低融点相が生成する場合と生成しない場合がある。しかし、この低融点が生成した場合は高温での溶体化熱処理を行うことができず、高温強度を向上させることができない。また、低融点相が生成しない条件で鋳造された試料での溶体化熱処理の予備検討結果を、同じ組成でも低融点相が生成する条件で鋳造された試料に適用した場合、低融点部が部分的に溶融して高温強度が大幅に低下する。以上の結果から、本発明では、Ｎｂの好適な添加量は０．３〜１％，より好ましくは０．６〜１．０％である。
【００５３】
ＴｉはＮｂと同様に硫化物を形成しやすく、燃料中にＳが多く存在する環境での耐食性を向上する効果がある。しかし、ＴｉもＮｂと同様に共晶部の融点を低下させるため、本発明ではＴｉの添加量は０．４％未満とした。尚、Ｔｉは不純物として含有される場合を除き特に添加するものではない。不純物として含有される量は０．００２〜０．００５％である。特に、０．０３％以下，より０．０１％以下が好ましい。
【００５４】
本発明合金のように結晶粒界の存在を許容して使用する場合には、Ｓｉ，Ｍｎ，Ｐ，Ｓ，Ｍｇ，Ｃａ等の不純物量を厳密に制限する必要がある。本発明では、これらの元素を故意には添加しなかったが、添加元素及びＮｉ中に不純物として存在し、合金中に混入する恐れがある。従って、本発明ではこれらの元素の最大値を下記のとおりに制限して鋳造した。
【００５５】
Ｓｉ≦０．０５％，Ｍｎ≦０．５％，Ｐ≦０．００５％，Ｓ≦０．００３％，Ｍｇ≦１００ｐｐｍ，Ｃａ≦１００ｐｐｍ
さらに、Ｆｅ及びＣｕも不純物レベルとすることが好ましく、本発明では共に０．２％以下が好ましい。また、合金中のガス量も［Ｎ］：１５ｐｐｍ未満，［Ｏ］：１５ｐｐｍ未満が好ましい。
【００５６】
本発明合金はＹあるいはＬａ，Ｃｅ等の希土類元素を添加することが可能である。これらの元素は耐酸化性の向上に有効であるが、鋳造時に鋳型材と反応して表面欠陥をつくりやすいこと、また合金の初期溶融温度を著しく低下させるため、本発明合金にこれらの元素を添加する場合は、これらの元素の総量を０．５％以下にすることが好ましい。
【００５７】
更に、表３に示すように、本発明合金は凝固方向でのクリープ破断時間が３５０時間以上のもの、更に５００時間以上のものが得られるとともに、その垂直方向でのクリープ破断時間は前者に対して３０時間、或いは更に後者に対して４５時間以上のものが得られることが分る。その結果、特に、凝固方向の１０４０℃，１４ｋｇ／ｍｍ^２でのクリープ破断時間が３５０時間以上で、それに垂直な方向での９２０℃，３２ｋｇ／ｍｍ^２でのクリープ破断時間が凝固方向のクリープ破断時間を０．１５倍した値より３２．５を差し引いた値以上とする横方向での高いクリープ破断強度が得られる。
【００５８】
〔実施例３〕
表４に示すＮｏ．３４の組成の多結晶マスターインゴットを真空誘導溶解で作製し、一方向凝固炉で１５ｍｍ×１００ｍｍ×２２０ｍｍの一方向凝固平板を鋳造した。この合金について、Ｎｏ．１〜２５と同様の方法による予備検討結果をもとにして、１２７５℃，１，４，２０時間の溶体化熱処理を施した試料及び時効熱処理のみの試料を用意した。
【００５９】
また、比較合金として、従来一方向凝固用合金のＣＭ１８６ＬＣを同時に評価した。まず、この合金の多結晶マスターインゴットを米国特許第５，０６９，８７３号公報中の表２に示される組成を目標にして真空誘導溶解で作製し、続いて一方向凝固炉で１５ｍｍ×１００ｍｍ×２２０ｍｍの一方向凝固平板を鋳造した。比較合金１はこの合金に米国特許第５，０６９，８７３号公報中に示された１０８０℃／４時間／空冷＋８７１ ℃／２０時間／空冷の熱処理を施したものである。さらに、この合金についてもＮｏ．１〜２５と同様の方法で最適溶体化熱処理温度を求めた。その結果、初期溶融温度が１２７７℃であることが明らかになったため、溶体化熱処理温度を１２７５℃とし、この温度で処理時間を１，４，８，２０及び４０時間とした比較合金２〜６も併せて評価した。
【００６０】
【表４】

【００６１】
なお、Ｎｏ．３４及び比較合金ともに溶体化熱処理後は空冷とし、これに続く時効熱処理の条件は１０８０℃／４時間／空冷＋８７１℃／２０時間／空冷とした。
【００６２】
以上の試料について、１０４０℃，１４ｋｇｆ／ｍｍ^２の条件で凝固方向のクリープ破断強度を、９２７℃，３２ｋｇｆ／ｍｍ^２の条件で横方向のクリープ破断強度を評価した。結果を表４に示す。また、画像解析で求めた体積率で表した溶体化率と凝固方向の強度及び横方向の強度の関係を図７に示す。この関係から、本発明合金３４は、比較合金より短い溶体化熱処理時間、つまり、少ない溶体化率で、比較合金より優れた凝固方向のクリープ強度、つまり、高温強度を示すことがわかる。従って、本発明合金は、横方向の強度、つまり結晶粒界の強度を低下させることなく高温強度を向上することが可能である。これは、本発明合金が比較合金と比べて著しく高い量のＴａを含有していることで、結晶粒界近傍のγ′相の固溶温度が結晶粒内と比べてかなり高くなり、結晶粒界近傍のγ′相をγ相中に固溶させることなく、結晶粒内のγ′相を溶体化することが可能となり、結晶粒界からＢを拡散消失させることなく結晶粒内の強度を向上させることができたためと考えられる。また、同等の溶体化率においても高温強度が優れているのは、相対的に低いＣｏ量の影響と考えられる。
【００６３】
〔実施例４〕
表３中のＮｏ．６１の組成をベースに１５０ｋｇマスターインゴットを作製した。分析結果（重量％）を表５に示す。残部はＮｉである。このマスターインゴットを用いて、溶解量約３．４ｋｇ，直径１５ｍｍ×高さ１８０ｍｍの丸棒８本取りのセレクター方式鋳型で単結晶丸棒試料を鋳造した。単結晶鋳造後、塩酸＋過酸化水素水を用いてマクロエッチングし、丸棒試料が単結晶となっていることを確認した。また、背面ラウエＸ線で結晶方位を測定し、丸棒試料の軸方向の結晶方位が〈００１〉方位から１０°以内の試料のみを選択した。この丸棒試料から直径６．３５ｍｍ，標点間距離２５．４ｍｍのクリープ歪み測定用ツバ付き試験片を採取し、単結晶試料のクリープ強度を測定した。結果を表６に示す。
【００６４】
【表５】

【００６５】
表６の結果をラーソン・ミラーパラメータで整理した結果を図８に示す。比較のために、Ｕ．Ｓ．Ｐ．５，３９９，３１３に示されるｌｏｗａｎｇｌｅｂｏｕｎｄａｒｉｅｓの強度を改善した単結晶合金のデータを示す。比較合金の単結晶の強度はＵ．Ｓ．Ｐ．５，３９９，３１３中で最も特性の優れたＮｏ．４９合金の学会発表に相当する“Ｒｅｎｅ’Ｎ４：Ａｆｉｒｓｔｇｅｎｅｒａｔｉｏｎｓｉｎｇｌｅｃｒｙｓｔａｌｔｕｒｂｉｎｅａｉｒｆｏｉｌａｌｌｏｙｗｉｔｈｉｍｐｒｏｖｅｄｏｘｉｄａｔｉｏｎｒｅｓｉｓｔａｎｃ，ｌｏｗａｎｇｌｅｂｏｕｎｄａｒｙｓｔｒｅｎｇｔｈａｎｄｓｕｐｅｒｉｏｒｌｏｎｇｔｉｍｅｒｕｐｔｕｒｅｓｔｒｅｎｇｔｈ”；Ｅ．Ｗ．ＲｏｓｓａｎｄＫ．Ｓ．Ｏ’Ｈａｒａ；Ｓｕｐｅｒａｌｌｏｙｓ１９９６，ＴＭＳ，（１９９６）ｐｐ１９−２５ｒｅｓｉｓｔａｎｃ，ｌｏｗａｎｇｌｅｂｏｕｎｄａｒｙｓｔｒｅｎｇｔｈａｎｄｓｕｐｅｒｉｏｒｌｏｎｇｔｉｍｅｒｕｐｔｕｒｅｓｔｒｅｎｇｔｈ”；Ｅ．Ｗ．ＲｏｓｓａｎｄＫ．Ｓ．Ｏ’Ｈａｒａ；Ｓｕｐｅｒａｌｌｏｙｓ１９９６，ＴＭＳ，（１９９６）ｐｐ１９−２５の図７から読みとった。また、比較合金のｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓのｔｒａｎｓｖｅｒｓｅｄｉｒｅｃｔｉｏｎのデータはＵ．Ｓ．Ｐ．５，３９９，３１３中の表４のＮｏ．４９合金のものを用いた。
【００６６】
【表６】

【００６７】
図８のｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓのｔｒａｎｓｖｅｒｓｅｄｉｒｅｃｔｉｏｎの強度の比較から、結晶粒界が有る場合の強度はＮｏ．６１が比較合金より著しく優れていることがわかる。また、単結晶の場合の強度もＮｏ．６１が比較合金より優れており、さらに、Ｎｏ．６１のｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓのｓｏｌｉｄｉｆｉｅｄｄｉｒｅｃｔｉｏｎの強度が比較合金の単結晶材の強度を上回っている。Ｎｏ．６１が比較合金より結晶粒界が有る場合の強度が優れている原因としては、結晶粒界元素Ｃ，Ｂの添加量が比較合金と比べて高いことが挙げられる。特に、結晶粒界の強度向上に最も効果のあるＢの添加量が著しく多いためと考えられる。通常、結晶粒界強度の向上を目的にＢの添加量を増やすと合金の融点が低下し、完全な溶体化が不可能となる。しかし、Ｎｏ．６１のように完全な溶体化をしなくても比較合金より単結晶あるいはｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓのｓｏｌｉｄｉｆｉｅｄｄｉｒｅｃｔｉｏｎの強度が優れているのは、Ｒｅの添加，高いＴａ量，低いＴｉ及びＣｏ量の効果と考えられる。特に、融点を低下させるＴｉはＮｏ．６１では実質的に無添加としている。
【００６８】
本発明合金はｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓとして用いることができるが、本発明合金により単結晶動翼あるいは静翼を鋳造した場合の効果として、以下のことが考えられる。まず、結晶粒界の方位差はＵ．Ｓ．Ｐ．５，３９９，３１３に示された合金では実質的に約１２°までしか許容できないが、本発明合金では実質的に方位差がランダムなｃｏｌｕｍｎａｒｇｒａｉｎｃａｓｔｉｎｇｓのレベルまで許容することが可能である。従って、特に大型単結晶動翼あるいは静翼の歩留まり及び信頼性を向上することが可能である。また単結晶で鋳造することで、弾性定数の小さい方位を特定の方向に合わせることが可能となり、熱応力を低減し、動翼及び静翼の寿命を延長する効果がある。さらに、本発明の合金のように完全に溶体化されなくても優れた高温強度を示す合金は、完全溶体化時に大きく成長する再結晶粒の成長を最低限に抑さえることが可能である。従って、単結晶合金の強度をほとんど０まで低下させてしまう再結晶の問題を解決することができる。
【００６９】
【発明の効果】
本発明は、鋳造時の凝固割れを防止し、さらに使用中の信頼性を確保するのに十分な結晶粒界の強度を有し、かつ優れた高温強度をあわせ持つ方向性凝固用高強度Ｎｉ基超合金にある。本発明合金をガスタービンの高温部材に適用することによりガスタービンの燃焼温度の向上及び発電用ガスタービンの発電効率の一層の向上が発揮される。
【図面の簡単な説明】
【図１】Ｃを約０．１％含み、Ｈｆ及びＺｒを実質的に無添加とした場合のＢ量と凝固方向及び凝固方向と垂直方向のクリープ破断強度との関係を示す図。
【図２】Ｂを約０．０１％含み、Ｈｆ及びＺｒを実質的に無添加とした場合のＣ量と凝固方向及び凝固方向と垂直方向のクリープ破断強度との関係を示す図。
【図３】Ｃ，Ｈｆ及びＺｒを実質的に無添加とした場合のＢ量と凝固方向及び凝固方向と垂直方向のクリープ破断強度との関係を示す図。
【図４】Ｃを約０．１％，Ｂを約０．０１％含み、Ｈｆを実質的に無添加とした場合のＺｒ量と凝固方向及び凝固方向と垂直方向のクリープ破断強度との関係を示す図。
【図５】Ｃを約０．１％，Ｂを約０．０１％含み、Ｚｒを実質的に無添加とした場合のＨｆ量と凝固方向及び凝固方向と垂直方向のクリープ破断強度との関係を示す図。
【図６】Ｃを約０．１％，Ｂを約０．０１％含み、Ｚｒを実質的に無添加とした場合のＨｆ量と凝固方向と垂直方向の高温引張特性との関係を示す図。
【図７】本発明合金と比較合金における、溶体化率と凝固方向及び横方向のクリープ破断強度との関係を示す図。
【図８】本発明合金と比較合金のパラメータＰと応力との関係を示す線図。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a novel Ni-base superalloy used for high-temperature components such as moving blades and stationary blades of high-temperature equipment such as gas turbines, and particularly has excellent high-temperature strength and is difficult to produce at high yield. The present invention relates to a Ni-based superalloy most suitable for a high-temperature member made of a single crystal alloy having a large complex shape.
[0002]
[Prior art]
2. Description of the Related Art The combustion gas temperature of a gas turbine tends to increase year by year for the purpose of improving thermal efficiency, and each high-temperature member of the gas turbine requires a material having higher high-temperature strength than before. Therefore, the material for the rotor blade, which is exposed to the harshest environment among the high-temperature components of a gas turbine, changes from a normal cast material of a Ni-base superalloy to a columnar crystal material, and further in a gas turbine for an aircraft engine, Single crystal materials having higher strength at higher temperatures have been put to practical use. Here, the columnar crystal material and the single crystal material are one type of directional solidification material, and both are cast by a method called a unidirectional solidification method. In the columnar crystal material, the crystal grains are elongated in one direction by a method described in US Pat. No. 3,260,505 and the like, and the grain boundaries perpendicular to the direction of the main stress are reduced as much as possible to increase the high temperature strength. Improvement is possible. Further, in the case of a single crystal material, the high temperature strength can be further improved by forming the entire casting into one crystal by the method described in US Pat. No. 3,494,709.
[0003]
Furthermore, in order to improve the high-temperature strength of the Ni-based superalloy, a solution heat treatment for uniformly and finely precipitating the γ ′ phase as a precipitation strengthening phase is effective. That is, these Ni-base superalloys are ₃ It is strengthened by the precipitation of a γ 'phase mainly composed of (Al, Ti, Nb, Ta), and it is preferable that the γ' phase is uniformly and finely precipitated. However, in the as-solidified state, there is a coarse γ 'phase (a γ' phase precipitated and coarsened during cooling after solidification and a eutectic γ 'phase coarsely crystallized in the final solidified portion). Improve high-temperature strength by heating to a high temperature to form a solid solution of the γ 'phase into the γ phase of the matrix, then quenching (solution heat treatment), and precipitating as a uniform and fine γ' phase during the subsequent aging treatment Becomes possible. This solution heat treatment is preferably performed at a temperature as high as possible above the solid solution temperature of the γ 'phase and below the initial melting temperature. This is because the higher the heat treatment temperature, the more the region where the γ ′ phase is uniformly refined, and the higher the region where the γ ′ phase is uniformly refined, the higher the high-temperature strength. Another reason that the high-temperature strength of the single crystal material is excellent is that the solution heat treatment temperature is reduced by using an alloy exclusively for single crystal that contains only a grain boundary strengthening element that significantly lowers the initial melting temperature to the extent of impurities. This is because it is possible to increase the height, and as a result, almost all of the γ ′ phase coarsely precipitated after solidification can be uniformly refined.
[0004]
As described above, a single crystal material of a Ni-base superalloy having the highest high-temperature strength in the current technology as a blade material for a gas turbine is a CMSX-4 (US Pat. No. 4,643, U.S. Pat. No. 782), PWA 1484 (U.S. Pat. No. 4,719,080) and Rene'N5 (Japanese Unexamined Patent Application Publication No. 5-59474) have been developed and used as a moving blade of a gas turbine for an aircraft engine. Has been applied. However, as described above, all of these single crystal alloys contain only grain boundary strengthening elements such as C, B and Hf at the impurity level. Therefore, if there is a grain boundary in a blade cast using a single crystal alloy, the strength is extremely reduced, and in some cases, a vertical crack already occurs along the grain boundary during solidification. Therefore, in order to use a rotor blade cast with a single crystal alloy in a gas turbine, it is necessary to make the entire rotor blade a complete single crystal. Since the entire length of a rotor blade of a gas turbine for an aircraft engine is at most about 100 mm, the probability of generation of crystal grain boundaries during casting is small, and a single crystal alloy can be produced with a certain yield. However, the rotor blade of the gas turbine for power generation has a total length of about 150 to 450 mm, and it is very difficult to make the entire rotor blade a complete single crystal. Therefore, it is difficult to produce a rotor blade of a gas turbine for power generation at a high yield using a single crystal alloy with the current technology.
[0005]
As described above, for the purpose of improving the high-temperature strength of large rotor blades to which a single-crystal alloy cannot be applied due to the casting yield, development of an alloy for columnar crystal materials having excellent high-temperature strength has been promoted. As a result, CM186LC (US Columnar crystal alloys such as Patent No. 5,069,873) and Ren'142 (US Patent No. 5,173,255) have been developed. These alloys contain a sufficient amount of grain boundary strengthening elements to prevent solidification cracking during casting, ensure reliability during use, and have a PWA 1480 (US Pat. No. 4,209,348). ), High temperature strength comparable to first-generation single crystal alloys such as CMSX-2 (U.S. Pat. No. 4,582,548) or Ren'N4 (U.S. Pat. No. 5,399,313). are doing. Therefore, by using these columnar crystal alloys, it has become possible to produce, at a higher yield, a moving blade having a high-temperature strength substantially equal to that of a first-generation single-crystal alloy moving blade. However, the combustion gas temperature of gas turbines tends to further increase, and these conventional columnar alloys have become insufficient in high-temperature strength to meet the demand for further improvement in thermal efficiency of gas turbines.
[0006]
Further, JP-A-7-145703 and JP-A-5-59473 disclose a single crystal alloy containing a columnar crystal containing C, B, Zr and Hf.
[0007]
In view of the above, in order to increase the efficiency of a gas turbine for power generation, it is essential to develop an alloy that achieves both high production yield and high-temperature strength, which are inconsistent with the prior art.
[0008]
[Problems to be solved by the invention]
As described above, in order to improve the high-temperature strength of the Ni-based superalloy, it is effective to raise the solution heat treatment temperature as much as possible. For that purpose, it is preferable that the addition amount of the grain boundary strengthening element is set to about the impurity. . However, on the other hand, in order to ensure a high production yield and reliability during use, it is necessary for the alloy to include a grain boundary strengthening element for imparting appropriate strength to the grain boundaries. Therefore, conventionally, it was necessary to sacrifice the strength of the crystal grain boundary in order to improve the high temperature strength, and conversely, it was necessary to sacrifice the high temperature strength in order to improve the strength of the crystal grain boundary.
[0009]
Further, according to a study conducted by the present inventors on a conventional columnar crystal alloy CM186LC, (Materials and Process, Vol. 7 (1994), p1797 and Materials and Processes, Vol. 8 (1995), p1458), It has been clarified that B, which is a grain boundary strengthening element, diffuses from crystal grain boundaries into grains during solution heat treatment. Therefore, although this alloy contains a grain boundary strengthening element, if solution heat treatment is performed to improve the high-temperature strength, the strength of the grain boundaries will be reduced to a level that cannot be practically used. Usually, in the directional solidification material, the direction in which the main stress acts is matched with the solidification direction, and thus the high-temperature strength is evaluated by the strength in the solidification direction. In this case, the high-temperature strength, that is, the strength in the solidification direction parallel to the crystal grain boundaries, increases as the solution of the γ 'phase progresses, but conversely, the strength of the crystal grain boundaries, that is, the solidification direction perpendicular to the crystal grain boundaries, increases. The vertical strength will be reduced.
[0010]
From the above, even if the grain boundary strengthening element is simply added to an existing single crystal alloy, an improvement in yield can be expected, but an excellent high-temperature strength cannot be expected because the solution heat treatment temperature is significantly lowered. Further, with respect to the conventional columnar crystal alloy, the solution heat treatment temperature cannot be further increased due to the problems of initial melting and reduction in the strength of crystal grain boundaries, and improvement in high-temperature strength beyond the current level cannot be expected.
[0011]
It is an object of the present invention to prevent solidification cracking during casting, to have sufficient strength at the grain boundaries to ensure reliability during use, and to combine high-temperature strength superior to conventional columnar alloys. An object of the present invention is to provide a high-strength Ni-base superalloy for directional solidification.
[0012]
[Means for Solving the Problems]
The present invention aims to obtain an alloy composition that achieves both high-temperature strength and crystal grain boundary strength, which were considered to be in a contradictory relationship in the prior art, and to obtain four types of crystals of C, B, Hf, and Zr. The results were obtained by adding the grain boundary strengthening elements to the single crystal alloy in various combinations, and examining the relationship between the amount of the grain boundary strengthening elements and the high temperature strength, the strength of the grain boundaries, and the solution heat treatment.
[0013]
A grain boundary strengthening element was added to a polycrystalline master ingot prepared as a component of a single crystal alloy in a unidirectional solidification furnace to cast a unidirectional solidified plate having a desired composition. Here, the high-temperature strength of each composition was evaluated by the creep rupture strength in the solidification direction. In addition, the strength of the grain boundaries to ensure castability and reliability during use is the creep rupture strength in the direction perpendicular to the solidification direction of the plate, that is, the direction in which the grain boundaries are perpendicular to the stress load direction. Evaluated by high temperature tensile properties.
[0014]
As a result, in an area considerably higher than the conventionally known optimum addition amount of B, there is an optimum addition amount that balances the strength in the solidification direction and the direction perpendicular to the solidification direction, that is, the high-temperature strength and the strength of the crystal grain boundary. It became clear. That is, C is added as a grain boundary strengthening element in an amount of 0.03 to 0.20%, preferably 0.05 to less than 0.1%, and Hf is 1.5% or less, preferably 0.01 to 0.5%. % And less than 0.02%, preferably less than 0.01%, the optimum B content for both the strength in the solidification direction and the strength in the direction perpendicular to the solidification direction is 0.004 to 0.05. %, Preferably more than 0.015% and not more than 0.04%, and it is apparent that both the strength in the solidification direction and the strength in the direction perpendicular to the solidification direction are maximized at around 0.03%. became. Preferably, the amount of B added to the conventional columnar crystal alloy is about 0.015%, which is about twice the amount of the conventional alloy.
[0015]
B is an element that significantly lowers the initial melting temperature of the alloy. Therefore, when a large amount of B is added, it is necessary to consider the effect of a decrease in the initial melting temperature. No decrease in melting temperature was observed.
[0016]
C is also an important element for achieving both high-temperature strength and crystal grain boundary strength. That is, in an alloy containing 0.007 to 0.015% of B, less than 0.5% of Hf, and less than 0.01% of Zr as a grain boundary strengthening element, the creep rupture strength in the solidification direction is determined by the amount of C added. It became clear that it decreased as it increased. On the other hand, the creep rupture strength in the direction perpendicular to the solidification direction increases with the addition amount up to 0.20% or less, preferably up to 0.10%. It was found that the peak was reduced. Therefore, if only the creep rupture strength in the direction perpendicular to the solidification direction is considered, the optimum amount of C added is around 0.1%, while the creep rupture strength in the solidification direction increases with the increase in C content. Taking into account that the C content is lowered, it is considered that the optimum range of the amount of C for satisfying both the high-temperature strength and the strength of the crystal grain boundary is 0.05 to less than 0.1%. When the addition amount of C is less than 0.03% or less than 0.05%, the high-temperature strength is excellent, but the strength of the grain boundaries is low, solidification cracking during casting is prevented, and reliability during use is reduced. Nature cannot be secured. On the other hand, when the addition amount of C is 0.2% or more, preferably 0.1% or more, the high temperature strength is remarkably reduced and the strength of the crystal grain boundaries is also reduced.
[0017]
Zr and Hf are homologous elements and are known to have substantially the same effect on Ni-based superalloys. Studies in the present invention have revealed that Zr significantly lowers the initial melting temperature of the alloy and makes it impossible to perform solution heat treatment at high temperatures, thereby lowering the creep rupture strength in the solidification direction of the alloy. Further, since it has been found that Zr has no effect on the creep rupture strength in the lateral direction, it is desirable that the amount of Zr added be less than 0.01%, and more preferably that Zr is substantially free of Zr. It must be added. Hf, like Zr, significantly lowers the initial melting temperature of the alloy and makes it impossible to perform solution heat treatment at a high temperature, thereby lowering the creep rupture strength in the solidification direction of the alloy and the creep rupture strength in the transverse direction. Has almost no effect. However, Hf has the effect of improving the tensile elongation in the transverse direction, and when added in an amount of about 0.25%, although the creep rupture strength in the solidification direction is slightly reduced, the creep rupture strength and the tensile strength in the transverse direction are reduced. It was clear that both improved. Therefore, the addition amount of Hf is desirably set to 0.01 to less than 0.5%, and a preferable range of the addition amount is 0.2 to 0.4%. Further, the range of the optimum amount of addition for satisfying both the high temperature strength and the strength of the crystal grain boundary is 0.2 to 0.3%.
[0018]
By optimizing the combination of the amounts of B, C, Hf and Zr as described above, the alloy of the present invention contains a sufficient grain boundary strengthening element to obtain the strength of the grain boundary, and By making it possible to perform a sufficient solution heat treatment to improve the strength, it has become possible to achieve both high-temperature strength and strength of crystal grain boundaries, which could not be achieved by the prior art.
[0019]
In addition, the above result is not determined only by the combination of the grain boundary strengthening elements, and the effect of the element contributing to strengthening in the crystal grains cannot be ignored. One of the features other than the grain boundary strengthening element in the composition of the alloy of the present invention is the amount of Co added. Many of the conventional columnar alloys contain a large amount of Co of 9% or more. However, in the present invention, the large amount of Co significantly reduces the high-temperature strength of the alloy and reduces the strength of the grain boundaries. Has little effect. However, on the other hand, Co has the effect of improving the corrosion resistance in a combustion gas atmosphere, and the material of the moving blade and the stationary blade for a gas turbine for power generation for which the corrosion resistance is emphasized has a high temperature strength within a range that does not significantly lower the high-temperature strength. Addition is essential.
[0020]
Further, in the alloy of the present invention, the high-temperature strength could be improved without lowering the strength of the crystal grain boundaries by the solution heat treatment, which is also a major factor because the amount of Ta added could be optimized. As described above, in the conventional columnar crystal alloy CM186LC, when solution heat treatment is performed to improve the high-temperature strength, B diffuses from the crystal grain boundaries into the crystal grains, and the strength of the crystal grain boundaries is extremely reduced. I do. This is because when the γ 'phase is once dissolved in the γ phase by the solution heat treatment, B diffuses into the γ phase almost simultaneously with the γ' phase near the crystal grain boundaries forming a solid solution in the γ phase. This is because B starts and eventually disappears from the crystal grain boundaries. To solve this problem, the alloy of the present invention added a much higher amount of Ta than the conventional alloy. As a result, the solid solution temperature of the γ ′ phase near the crystal grain boundary becomes considerably higher than that inside the crystal grain, and the γ ′ phase near the crystal grain boundary does not form a solid solution in the γ phase. 'Solution can be dissolved. Therefore, in the alloy of the present invention, the strength in the crystal grain can be improved without diffusing and eliminating B from the crystal grain boundary, and as a result, the high-temperature strength can be improved without lowering the strength of the crystal grain boundary. Became possible.
[0021]
Generally, the higher the solution solution rate, that is, the higher the proportion of the region in which the γ ′ phase is refined in the alloy, the higher the high-temperature strength. However, considering the strength of the crystal grain boundaries as described above, the solution solution rate is as low as possible. Preferably, it is small. Therefore, in order to achieve both the high-temperature strength of the alloy and the strength of the crystal grain boundaries, a composition that can exhibit excellent high-temperature strength even when the solution solution rate is small is preferable. Therefore, in the alloy of the present invention, the amounts of Re and W effective for solid solution strengthening are optimized, and the alloy is solid-solution strengthened as much as possible, thereby enabling high-temperature strength to be improved at a relatively low solution rate. .
[0022]
The alloy of the present invention is suitable for directional solidification by a unidirectional solidification method. In particular, in the case of a moving blade for a gas turbine, it is desirable that casting be performed with the direction in which centrifugal force acts as a solidification direction. Further, although the description has been made on the premise that it is used for a moving blade for a gas turbine, it can also be used for other high-temperature components such as a stationary blade for a gas turbine. In the case of a gas turbine stationary blade, it is preferable to use the solidification direction in a direction in which the thermal stress is maximized. The alloy of the present invention can be used not only for ordinary columnar crystal moving blades and columnar crystal stationary blades but also for moving blades in which a grain boundary is generated in a part of the moving blade during single crystal casting. is there. Conventionally, such a moving blade was regarded as a defective product, but if the alloy of the present invention is used, such a moving blade can be sufficiently used, and as a result, the casting yield of a single crystal moving blade can be significantly improved. Becomes possible. Further, the alloy of the present invention can be used for a normal single crystal blade. Even if it is possible to cast complete single-crystal blades and single-crystal stationary blades with a high yield using conventional single-crystal alloys, the alloy of the present invention can simplify the inspection for the presence or absence of grain boundaries. It is possible to reduce production costs. Furthermore, although the presence or absence of crystal grain boundaries on the inner surface of the rotor blade was conventionally guaranteed by a sampling fracture test, the alloy of the present invention can guarantee the strength even if there is a crystal boundary, greatly improving the reliability of the rotor blade. It is possible to improve.
[0023]
As described above, in the present invention, C: 0.03 to 0.20%, preferably 0.05% to less than 0.1%, and B: 0.004 to 0.05%, preferably 0.015% Over 0.04%, Hf: 1.5% or less, preferably 0.01 to less than 0.5%, Zr: 0.02% or less, preferably less than 0.01%, Cr: 1.5 to 16% %, Mo: 6% or less, W: 2 to 12%, Re: 0.1 to 9%, Ta: 2 to 12%, Nb: 4.0 or less, preferably 0.3 to 4%, Al: 4 0.0-6.5%, Ti: Less than 0.4%, preferably not added, Co: High-strength Ni-base for directional solidification with excellent grain boundary strength containing 9% or less and 60% or more Ni It is a superalloy. In particular, alloys exhibiting both high-temperature strength and crystal grain boundary strength at the same time and exhibiting excellent corrosion resistance in a combustion gas are C: 0.05 to 0.10% and B: 0.018 to 0% by weight. 0.04%, Hf: 0.01 to less than 0.5%, Zr: less than 0.01%, Cr: 4 to 12.5%, Mo: 4.5% or less, W: 5 to 10%, Re: 1 to 6%, Ta: 5 to 12%, Nb: 0.3 to 3%, Al: 4.0 to 6.0%, Co: 0.5 to 1.2%, and the remaining unavoidable impurities and Ni This is a high-strength Ni-based superalloy for directional solidification having excellent crystal grain boundary strength. Further, when an alloy having higher high-temperature strength is required, C: 0.06 to 0.10%, B: 0.018 to 0.035%, Hf: 0.1 to 0% by weight. 0.5%, Cr: 6.5 to 8.5%, Mo: 0.4 to 3.0%, W: 5.5 to 9.5%, Re: 1.0 to 6.0%, Ta: 6 to 10.5%, Nb: 0.3 to 1.55%, Al: 4.0 to 6.0%, Co: 0.5 to 2.5%, and crystal grains comprising Ni and unavoidable impurities. A high-strength Ni-base superalloy for directional solidification having excellent field strength is suitable, and more preferable compositions are C: 0.06 to 0.10% and B: 0.018 to 0.035 by weight%. %, Hf: 0.2-0.3%, Cr: 6.9-7.3%, Mo: 0.7-2.0%, W: 7.0-9.0%, Re: 1. 2 to 2.0%, Ta: 8.5 to 9.5%, Nb 0.6-1%, Al: 4.0-6.0%, Co: 0.5-1.2%, Ti less than 0.4%, preferably not added and 60% or more Ni, preferably balance The composition is composed of unavoidable impurities and Ni.
[0024]
Particularly, in an environment in which the fuel contains a large amount of impurities such as S, C: 0.06 to 0.08%, B: 0.018 to 0.035%, and Hf: 0.2 to 0. 3%, Cr: 6.9 to 7.3%, Mo: 0.7 to 1%, W: 8 to 9%, Re: 1.2 to 1.6%, Ta: 8.5 to 9.5. %, Nb: 0.3 to 1%, Ti: less than 0.5%, Al: 4.9 to 5.2%, Co: 0.8 to 1.2%, and the balance consisting of unavoidable impurities and Ni A high-strength Ni-base superalloy for directional solidification, which is excellent in crystal grain boundary strength and characterized by:
[0025]
When a Ni-base superalloy having the above composition is used, the creep rupture life in the solidification direction is 1040 ° C. and 14 kgf / mm. ² Creep rupture life of 927 ° C, 32 kgf / mm in the direction perpendicular to the solidification direction for 350 hours or more under the conditions of ² It is possible to obtain a directionally solidified casting excellent in both high-temperature strength and crystal grain boundary strength for 30 hours or more under the above conditions.
When a Ni-base superalloy having the above composition is used, the γ 'phase can be formed into a rectangular parallelepiped having a side of 0.5 μm or less in a region having a volume ratio of 50% or more by solution heat treatment and solidification. Creep rupture life in the direction perpendicular to the direction is 927 ° C, 32 kgf / mm ² The tensile strength at 800 ° C. in the direction perpendicular to the solidification direction is 95 kgf / mm for 30 hours or more under the following conditions. ² A directionally solidified casting having the above characteristics and excellent in both high-temperature strength and crystal grain boundary strength can be obtained.
[0026]
In the present invention, C is 0.03 to 0.20%, B is 0.004 to 0.05%, Cr is 4.0 to 12.5%, Mo is 4.5% or less, W is 5.0 to 10.0%, Re 1.0 to 7.0%, Ta 5.0 to 12.0%, Nb 0.3 to 4.0%, Al 4.0 to 6.5%, Ti less than 0.4%, Co 0.5 to 5.0 %, Hf 1.5% or less, Zr 0.15% or less and 60% or more of Ni, and the C amount is not less than 0.15 minus a value obtained by multiplying the B amount by 5.45. Ni-based superalloy for directional solidification.
[0027]
In particular, the C amount and the B amount are not more than a straight line connecting (0.20%, 0.03%) and (0.08%, 0.05%), and more preferably (0.20%, 0.01%). %) And (0%, 0.047%).
[0028]
In the present invention, the solidification direction is 1040 ° C., 14 kg / mm ² Rupture time at 350 ° C or more and 920 ° C, 32 kg / mm in the direction perpendicular to the solidification direction ² The columnar crystal Ni-base superalloy has a creep rupture time of 30 hours or more. Particularly, the former is preferably 500 hours or more and the latter 45 hours or more.
[0029]
In the present invention, the solidification direction is 1040 ° C., 14 kg / mm ² Rupture time at 350 ° C or more and 920 ° C, 32 kg / mm in the direction perpendicular to the solidification direction ² Wherein the creep rupture time of the columnar Ni-based superalloy is not less than a value obtained by subtracting 32.5 from a value obtained by multiplying the creep rupture time in the solidification direction by 0.15. In particular, the former is preferably 500 hours or more.
[0030]
Further, in the present invention, the ratio of the Co amount to the Mo amount is preferably set to 0.2 to 5, more preferably 0.4 to 2.0.
[0031]
Table 1 shows a wide range, a desirable range, a preferable range, and an optimum range for each alloy composition according to the present invention.
[0032]
The above-mentioned Ni-based alloy according to the present invention is preferably such that the matrix γ phase is composed of a single crystal.
[0033]
[Table 1]

[0034]
BEST MODE FOR CARRYING OUT THE INVENTION
[Example 1]
Table 2 shows the relationship between the crystal grain boundary strengthening elements of C, B, Hf, and Zr, and the relationship between the added amount of the crystal grain boundary strengthening elements and the high-temperature strength and the strength of the crystal grain boundaries. Here, the base alloy is composed of 7.8Cr-7.2W-1.8Mo-4.7Al-1.6Nb-7.5Ta-1.6Re-balance Ni by weight%. These alloys are obtained by adding a grain boundary strengthening element in a unidirectional solidification furnace to a base alloy polycrystalline master ingot produced by vacuum induction melting, and forming a plate having unidirectionally solidified columnar crystals of 15 mm × 100 mm × 100 mm. Was cast. Next, several 10 mm × 10 mm × 10 mm blocks were collected from the unidirectionally solidified flat plate, and heat-treated at 1250, 1260, 1270, 1280, 1290, 1300, 1310, 1320, and 1330 ° C. for 2 hours. The optimal solution heat treatment conditions for each composition were determined from the subsequent structure. In addition, about some alloys, the temperature was changed in a finer range as needed, and more accurate optimal solution treatment conditions were examined. Here, the optimal heat treatment conditions are such that the γ 'phase can be arranged in a rectangular parallelepiped shape with one side of 0.5 μm or less in a region of 50% or more by volume, and as high as possible but not higher than the initial melting temperature. Table 2 shows the conditions of the solution heat treatment actually determined for each alloy determined from the preliminary study. Air cooling was performed after the solution heat treatment, and the conditions of the subsequent aging heat treatment were constant at 1080 ° C./4 hours / air cooling + 871 ° C./20 hours / air cooling for all alloys.
[0035]
Here, the high-temperature strength is determined at 920 ° C. and 32 kgf / mm for a sample collected in the solidification direction of the unidirectional solidified plate. ² The creep rupture strength under the following conditions was evaluated. Hereinafter, this is referred to as creep rupture strength in the solidification direction. The strength of the grain boundaries is perpendicular to the solidification direction of the unidirectionally solidified plate (hereinafter referred to as the transverse direction), that is, a sample collected so that the stress axis is perpendicular to the grain boundaries. 920 ℃, 32kgf / mm ² And the high temperature tensile properties at 800 ° C. These results are also shown in Table 2. The shape of the test piece was 6 mm in diameter and 30 mm in distance between gauge marks in both the creep rupture test and the high-temperature tensile test. These specimens appear to exhibit the same properties in the same direction as the solidification direction as if they were entirely single crystals.
[0036]
The width of the crystal grains of the unidirectionally solidified flat plate is about 1 to 5 mm on the solidification start side (lower part) and about 5 to 10 mm on the upper part. A sample for evaluating the strength in the lateral direction was collected at the center of the flat plate (the width of the crystal grains was about 5 mm). Therefore, about five crystal grain boundaries exist in the gauge length. No particular consideration was given to the sampling position in the solidification direction. In an extreme case, the inside of the gauge length may be a single crystal, but usually there are about three crystal grain boundaries.
[0037]
[Table 2]

[0038]
FIG. 1 shows the relationship between the amount of B and the creep rupture strength in the solidification direction and in the transverse direction when C is contained at about 0.1% and Hf and Zr are substantially not added. In this case, the optimum addition amount is around 0.03% in both the solidification direction and the lateral direction. Considering that the amount of B added to the conventional columnar crystal alloy is at a level of about 0.015%, this result shows that the amount of B added is about twice the amount of B that has been conventionally considered to be optimal. This indicates that the optimal value of When the B content is 0.017 to 0.040%, high strength can be obtained.
[0039]
FIG. 2 shows the relationship between the C content and the creep rupture strength in the solidification direction and in the transverse direction when B is contained at about 0.01% and Hf and Zr are substantially not added. FIG. 3 shows the relationship between the amount of B and the creep rupture strength in the solidification direction and the lateral direction when C is substantially not added in addition to Hf and Zr. From these relations, it can be seen that the creep rupture strength in the solidification direction decreases with the addition of C, but C is an essential element for obtaining strength in the transverse direction. Therefore, it is necessary to precisely control the amount of carbon in order to achieve both high-temperature strength and crystal grain boundary strength. Further, by controlling the C content, it is possible to obtain an alloy in which either high-temperature strength or crystal grain boundary strength is emphasized. Specifically, when the high-temperature strength is emphasized, the lower the C amount, the better. When the importance of the crystal grain boundary is more important than the high-temperature strength, the higher the C amount, the better.
[0040]
4 and 5 show the relationship between the amounts of Zr and Hf and the creep rupture strength in the solidification direction and the lateral direction when C is contained at about 0.1% and B is contained at about 0.01%. From these results, it became clear that Zr and Hf reduced the creep rupture strength in the solidification direction and hardly improved the creep rupture strength in the transverse direction. However, Hf has the effect of improving the tensile ductility in the transverse direction as shown in FIG.
[0041]
[Example 2]
A polycrystalline master ingot of an alloy having the composition shown in Table 3 was produced by vacuum induction melting, and a unidirectional solidified flat plate of 15 mm × 100 mm × 220 mm was cast in a unidirectional solidification furnace. With respect to these alloys, Similar to 1 to 25, several blocks of 10 mm × 10 mm × 10 mm are collected from the unidirectionally solidified flat plate, and each of the blocks is heated at 1250, 1260, 1270, 1280, 1290, 1300, 1310, 1320, 1330 ° C. for 2 hours. Preliminary examination of solution heat treatment conditions was conducted. Based on these results, the heat treatment temperature of these alloys was increased in steps of 1250 ° C./4 hours in increments of 10 ° C./4 hours and finally heat treated at the maximum solution heat treatment temperature shown in Table 2 for 4 hours. A post-air cooling multi-stage solution heat treatment was performed. The aging heat treatment following the solution heat treatment was 1080 ° C./4 hours / air cooling + 871 ° C./20 hours / air cooling for all alloys.
[0042]
Table 3 also shows the properties evaluation results of each alloy. Of these tests, the samples for the creep rupture test and the high temperature tensile test were No. Samples were collected in the same manner as in Nos. 1 to 25, and the shape of the test piece was also the same. The conditions were a creep rupture test in the solidification direction at a temperature of 1040 ° C. and a stress of 14 kgf / mm. ² , Transverse creep rupture test at a temperature of 927 ° C and a stress of 32 kgf / mm ² The temperature in the transverse tensile test was 800 ° C.
[0043]
Further, the corrosion resistance of some alloys was evaluated by a simulated combustion test on a real machine. The test piece was a round bar having a diameter of 9 mm and a length of 50 mm. In a simulated combustion gas atmosphere of an actual machine, corrosion loss after descaling after 900 times of cooling at 7 hours at 900 ° C. for 7 hours was measured.
[0044]
[Table 3]

[0045]
As shown in the table, it is desirable to add Ta at least 2% or more, preferably 5% or more in order to improve high-temperature strength, and the optimum value for high-temperature strength is in the range of 8.5 to 9.5%. Exists. On the other hand, a large amount of Ta increases the solid solution temperature of the γ 'phase as described above. Therefore, when Ta is added excessively, the difference between the initial melting temperature of the alloy and the solid solution temperature of the γ 'phase is reduced, and the region in which the γ' phase can be solutionized without initial melting is reduced, thereby strengthening the precipitation of the alloy. The amount decreases. Therefore, the addition of Ta exceeding 12% has no effect on improving the high-temperature strength anymore, and the upper limit is preferably set to 10% or less.
[0046]
No. Co has substantially the same addition amounts of the other alloy elements, and differs only in the Co addition amount. No. 38 and No. 38. From the relationship of 100 to 104, it is clear that the high-temperature strength decreases as the added amount increases. Therefore, in consideration of high-temperature strength, the amount of Co added is 9% or less, preferably less than 99%, and more preferably 0.5 to 5%. In particular, the addition of about 0.5 to 1.2% of Co has the effect of improving the corrosion resistance.
[0047]
W and Re are elements effective for improving the high-temperature strength by solid-solution strengthening the alloy, and it is desirable to add 2%, preferably 5% and 0.1%, preferably 1% or more, respectively. Further, when high-temperature strength is emphasized, it is preferable to add 5.5% and 1.2% or more, respectively. On the other hand, the effect of these elements saturates at a certain amount of addition, and if added excessively, the high-temperature strength of the alloy is rather lowered. This is because if these elements are excessively added beyond the solid solubility limit, needle-like or plate-like precipitates mainly composed of W and Re are deposited. Therefore, it is desirable that the upper limit of the added amount of W and Re is 12%, preferably 10% and 9%, and more preferably 6%. Further, in order to suppress the precipitation of a large amount of the precipitates, it is preferable that the amounts of W and Re added are 9.5% and 3.1% or less, respectively. The optimum addition amount of W in the alloy of the present invention is 8.0 to 9.0%, and the optimum addition amount of Re is 1.2 to 1.6%. And W is preferably 5 to 10%, more preferably 5.5 to 9.5%, and Re is preferably 1 to 6%, more preferably 1.2 to 1.3%.
[0048]
It is preferable to consider the optimum addition amount of W and Re by the total amount of W and Re, and the high temperature strength becomes maximum when the W + Re amount is in the range of 9.5 to 12%. On the other hand, if the amount of W + Re is less than 9.5%, the solid solution strengthening becomes insufficient and the high-temperature strength decreases. If the amount of W + Re exceeds 12%, the above precipitates precipitate in large quantities, and the creep strength particularly at 1000 ° C. or higher is greatly reduced.
[0049]
Al is an essential element for forming a γ 'phase which is a strengthening factor of the Ni-based superalloy. In addition, Al ₂ O ₃ The formation of the coating contributes to the improvement of oxidation resistance and corrosion resistance. Therefore, it is desirable that the addition amount of Al is at least 4.0% or more, preferably 4.5% or more. However, an excessive addition exceeding 6.5% increases the amount of the eutectic γ 'phase in the alloy. The alloy of the present invention is considered so as to be able to exhibit excellent high-temperature strength even without complete solution treatment by optimizing the addition amount of an element effective for solid solution strengthening. Therefore, it has excellent high-temperature strength even in the presence of the eutectic γ 'phase. However, in the case of creep damage, the amount of the eutectic γ 'phase is preferably small, because the eutectic γ' phase eventually becomes a starting point of a crack and accelerates the fracture time of the material. Therefore, the addition amount of Al is preferably 6.5% or less, more preferably 5.7% or less. In particular, 4.7 to 5.4%, more preferably 4.9 to 5.2%.
[0050]
Cr is Cr ₂ O ₃ Since there is an effect of forming a coating to improve the corrosion resistance and oxidation resistance of the alloy, it is desirable to add at least 1.5% or more, more preferably 4% or more. However, excessive addition of Cr promotes the precipitation of the above-mentioned precipitates composed of W and Re, so that it is necessary to reduce the amount of W or Re added which is effective for high-temperature strength. Therefore, when importance is placed on high-temperature strength, the upper limit of the amount of Cr added is desirably 16%, more preferably 12.5%. In particular, 6.5 to 8.5%, more preferably 6.9 to 7.3%.
[0051]
Mo has the same effect as W or Re. However, since the corrosion resistance of the alloy in a combustion gas atmosphere is remarkably reduced, when the corrosion resistance is emphasized, the addition amount of Mo is reduced to 6% or less, more preferably to 4.5% or less. It is desirable to limit. Further, when importance is attached to corrosion resistance, it is preferable that the amount of Mo added be 0.4 to 1%, more preferably 0.7 to 1%.
[0052]
Nb is an element of the same family as Ta and has almost the same effect on high-temperature strength, and is contained in an amount of 0.3 to 4%. Further, since sulfides are easily formed, in an environment in which S is present in a large amount in the fuel, there is an effect that the penetration of S into the alloy is delayed and the corrosion resistance is improved. However, in the present invention, when a certain amount or more of Nb and B is present in the alloy, a low melting phase containing Nb and B as a main component is formed in the eutectic portion to significantly lower the initial melting temperature of the alloy. Was revealed. The low melting point phase is generated by segregation during solidification, and may or may not be formed depending on casting conditions. However, when this low melting point is generated, the solution heat treatment at a high temperature cannot be performed, and the high-temperature strength cannot be improved. In addition, when applying the preliminary results of solution heat treatment on a sample cast under the condition that a low melting point phase is not generated to a sample cast under the condition that a low melting point phase is generated even with the same composition, the low melting point part Melting at high temperature and the high temperature strength is greatly reduced. From the above results, in the present invention, the suitable addition amount of Nb is 0.3-1%, more preferably 0.6-1.0%.
[0053]
Ti easily forms sulfides like Nb, and has an effect of improving corrosion resistance in an environment where a large amount of S is present in the fuel. However, since Ti also lowers the melting point of the eutectic portion similarly to Nb, the addition amount of Ti is set to less than 0.4% in the present invention. Note that Ti is not particularly added unless it is contained as an impurity. The amount contained as an impurity is 0.002 to 0.005%. In particular, it is preferably 0.03% or less, more preferably 0.01% or less.
[0054]
When the alloy of the present invention is used while allowing the existence of crystal boundaries, it is necessary to strictly limit the amount of impurities such as Si, Mn, P, S, Mg, and Ca. In the present invention, these elements are not intentionally added, but may be present as impurities in the added elements and Ni and may be mixed into the alloy. Therefore, in the present invention, the casting was performed with the maximum values of these elements limited as follows.
[0055]
Si ≦ 0.05%, Mn ≦ 0.5%, P ≦ 0.005%, S ≦ 0.003%, Mg ≦ 100 ppm, Ca ≦ 100 ppm
Further, it is preferable that Fe and Cu are also at the impurity level, and in the present invention, both are preferably 0.2% or less. Also, the gas amount in the alloy is preferably [N]: less than 15 ppm and [O]: less than 15 ppm.
[0056]
The alloy of the present invention can be added with Y or a rare earth element such as La or Ce. Although these elements are effective in improving oxidation resistance, they are liable to react with the casting material during casting to form surface defects, and significantly lower the initial melting temperature of the alloy. When added, the total amount of these elements is preferably set to 0.5% or less.
[0057]
Further, as shown in Table 3, the alloy of the present invention has a creep rupture time in the solidification direction of 350 hours or more, and further has a creep rupture time of 500 hours or more, and the creep rupture time in the vertical direction is larger than that of the former. 30 hours, or more than 45 hours for the latter. As a result, especially at 1040 ° C. and 14 kg / mm in the solidification direction ² Rupture time of 350 hours or more at 920 ° C., 32 kg / mm in a direction perpendicular thereto ² , A high creep rupture strength in the transverse direction is obtained in which the creep rupture time in the solidification direction is 0.15 times the value obtained by subtracting 32.5 from the value obtained by subtracting 32.5.
[0058]
[Example 3]
No. shown in Table 4. A polycrystalline master ingot having a composition of 34 was produced by vacuum induction melting, and a unidirectionally solidified flat plate of 15 mm × 100 mm × 220 mm was cast in a unidirectional solidification furnace. With respect to this alloy, On the basis of the results of the preliminary examination by the same method as in Examples 1 to 25, a sample subjected to solution heat treatment at 1,275 ° C. for 1, 4, and 20 hours and a sample only subjected to aging heat treatment were prepared.
[0059]
As a comparative alloy, a conventional alloy for unidirectional solidification, CM186LC, was simultaneously evaluated. First, a polycrystalline master ingot of this alloy was prepared by vacuum induction melting with the aim of the composition shown in Table 2 in U.S. Pat. No. 5,069,873, followed by a unidirectional solidification furnace of 15 mm × 100 mm × A 220 mm unidirectionally solidified plate was cast. Comparative alloy 1 was obtained by subjecting this alloy to a heat treatment of 1080 ° C./4 hours / air cooling + 871 ° C./20 hours / air cooling shown in US Pat. No. 5,069,873. Further, with respect to this alloy, The optimum solution heat treatment temperature was determined in the same manner as in Examples 1 to 25. As a result, it was found that the initial melting temperature was 1277 ° C., so that the solution heat treatment temperature was set to 1275 ° C., and the comparative alloys 2 to 6 were treated at this temperature for 1, 4, 8, 20, and 40 hours. Was also evaluated.
[0060]
[Table 4]

[0061]
In addition, No. Both the alloy No. 34 and the comparative alloy were air-cooled after the solution heat treatment, and the conditions of the subsequent aging heat treatment were 1080 ° C./4 hours / air cooling + 871 ° C./20 hours / air cooling.
[0062]
For the above sample, 1040 ° C, 14 kgf / mm ² Creep rupture strength in the solidification direction at 927 ° C, 32 kgf / mm ² The creep rupture strength in the transverse direction was evaluated under the following conditions. Table 4 shows the results. FIG. 7 shows the relationship between the solutionization ratio expressed by the volume ratio obtained by the image analysis and the strength in the solidification direction and the strength in the lateral direction. From this relationship, it can be seen that the alloy 34 of the present invention exhibits a better creep strength in the solidification direction, that is, a high-temperature strength than the comparative alloy, with a shorter solution heat treatment time than the comparative alloy, that is, with a lower solution rate. Therefore, the alloy of the present invention can improve the high-temperature strength without lowering the strength in the lateral direction, that is, the strength of the crystal grain boundaries. This is because the alloy of the present invention contains a much higher amount of Ta than the comparative alloy, so that the solid solution temperature of the γ 'phase near the crystal grain boundary becomes considerably higher than that in the crystal grain, It is possible to form a solution of the γ ′ phase in the crystal grains without dissolving the γ ′ phase in the vicinity of the boundary in the γ phase, and to increase the strength in the crystal grains without diffusing and eliminating B from the crystal grain boundaries. It is thought that it could be improved. Further, it is considered that the high temperature strength is excellent even at the same solution solution ratio because of the relatively low Co content.
[0063]
[Example 4]
No. in Table 3. A 150 kg master ingot was produced based on the composition of No. 61. Table 5 shows the analysis results (% by weight). The balance is Ni. Using this master ingot, a single crystal round bar sample was cast in a selector type mold having a dissolution amount of about 3.4 kg, a diameter of 15 mm and a height of 180 mm, and eight round bars. After the single crystal casting, macro etching was performed using hydrochloric acid + hydrogen peroxide solution, and it was confirmed that the round bar sample was a single crystal. Further, the crystal orientation was measured by back Laue X-rays, and only those samples in which the crystal orientation in the axial direction of the round bar sample was within 10 ° from the <001> orientation were selected. A test piece with a brim for measuring creep strain having a diameter of 6.35 mm and a gauge length of 25.4 mm was collected from the round bar sample, and the creep strength of the single crystal sample was measured. Table 6 shows the results.
[0064]
[Table 5]

[0065]
FIG. 8 shows the results obtained by rearranging the results in Table 6 using Larson-Miller parameters. For comparison, U.S. Pat. S. P. 5 shows data of a single crystal alloy with improved strength of low angle bonds shown in US Pat. No. 5,399,313. The strength of the single crystal of the comparative alloy is U.S.A. S. P. No. 5,399,313, which has the most excellent characteristics. "Rene'N4: A first generation single crystal turbine airfoil alloy with improved oxidation recommendation, lower angle reinforcement. W. Ross and K. S. O'Hara; Superalloys 1996, TMS, (1996) pp19-25 resistance, low angle boundary strength and superior long timerupt. -25 of Fig. 7. In addition, the data of "transverse direction" of the column grain castings of the comparative alloy was that of the No. 49 alloy of Table 4 in U.S. Pat.
[0066]
[Table 6]

[0067]
From the comparison of the intensity of the transversal direction of the column grain castings in FIG. It can be seen that 61 is significantly superior to the comparative alloy. In addition, the strength in the case of a single crystal is no. No. 61 is superior to the comparative alloy. The strength of the solidified direction of 61 column grain castings exceeds that of the single crystal material of the comparative alloy. No. The reason why the strength of the alloy 61 having a grain boundary is superior to that of the comparative alloy is that the addition amount of the grain boundary elements C and B is higher than that of the comparative alloy. In particular, it is considered that the addition amount of B, which is most effective in improving the strength of the crystal grain boundaries, is extremely large. Normally, when the addition amount of B is increased for the purpose of improving the crystal grain boundary strength, the melting point of the alloy decreases, and complete solution formation becomes impossible. However, no. The strength of the single crystal or columnar grain castings of the modified direction is superior to that of the comparative alloy without complete solution solution as in the case of 61, because of the effects of the addition of Re, the high Ta amount, the low Ti and Co amounts. Conceivable. In particular, Ti which lowers the melting point is No. 1; In No. 61, it is substantially not added.
[0068]
The alloy of the present invention can be used as column grain castings. The following effects can be considered when the alloy of the present invention casts a single crystal moving blade or a stationary blade. First, the misorientation of the crystal grain boundaries is described in U.S. Pat. S. P. No. 5,399,313, the alloys of the present invention can tolerate substantially only up to about 12 °, while the alloys of the present invention can tolerate substantially misaligned levels of columnar grain castings. Therefore, it is possible to improve the yield and reliability of a large single crystal blade or a stationary blade, in particular. In addition, by casting with a single crystal, it is possible to adjust the orientation having a small elastic constant to a specific direction, thereby reducing the thermal stress and extending the life of the moving blade and the stationary blade. Furthermore, an alloy that exhibits excellent high-temperature strength even when not completely solution-solutioned, such as the alloy of the present invention, can minimize the growth of recrystallized grains that grow greatly during complete solution-solution. Therefore, the problem of recrystallization that reduces the strength of the single crystal alloy to almost zero can be solved.
[0069]
【The invention's effect】
The present invention provides high-strength Ni for directional solidification, which has sufficient strength at grain boundaries to prevent solidification cracking during casting and further ensure reliability during use, and also has excellent high-temperature strength. In base superalloy. By applying the alloy of the present invention to a high-temperature member of a gas turbine, the combustion temperature of the gas turbine is improved, and the power generation efficiency of the gas turbine for power generation is further improved.
[Brief description of the drawings]
FIG. 1 is a graph showing the relationship between the amount of B, the solidification direction, and the creep rupture strength in the direction perpendicular to the solidification direction when C is contained at about 0.1% and Hf and Zr are substantially not added.
FIG. 2 is a diagram showing the relationship between the amount of C, the solidification direction, and the creep rupture strength in the direction perpendicular to the solidification direction when B is contained in about 0.01% and Hf and Zr are substantially not added.
FIG. 3 is a graph showing the relationship between the amount of B, the solidification direction, and the creep rupture strength in a direction perpendicular to the solidification direction when C, Hf and Zr are substantially not added.
FIG. 4 shows the relationship between the amount of Zr, the solidification direction, and the creep rupture strength in the direction perpendicular to the solidification direction when C is contained in about 0.1% and B is contained in about 0.01% and Hf is substantially not added. FIG.
FIG. 5 shows the relationship between the amount of Hf, the solidification direction, and the creep rupture strength in the solidification direction and the perpendicular direction when Zr is substantially not added, containing about 0.1% of C and about 0.01% of B. FIG.
FIG. 6 is a graph showing the relationship between the amount of Hf and the high-temperature tensile properties in the solidification direction and the perpendicular direction when Zr is contained substantially without Cr containing about 0.1% of C and about 0.01% of B. .
FIG. 7 is a diagram showing a relationship between a solution heat treatment rate and a creep rupture strength in a solidification direction and a lateral direction in an alloy of the present invention and a comparative alloy.
FIG. 8 is a diagram showing a relationship between stress and a parameter P of the alloy of the present invention and a comparative alloy.

Claims

By weight, C: 0.03% to 0.20%, B: 0.004% to 0.05%, Hf: 0.01% to 1.5%, Zr: 0% to 0.02. %, Cr: 1.5% to 16%, Mo: 0.4% to 6%, W: 2% to 12%, Re: 0.1% to 9%, Ta: 2% or more 12% or less, Nb: 0.3% to 4%, Al: 4% to 6.5%, Ti: 0% to less than 0.4%, Co: 0.5% to 9% and the balance A high-strength Ni-base superalloy characterized by being substantially composed of Ni.

By weight, C: 0.06% to 0.09%, B: 0.016% to 0.04%, Hf: 0.01% to less than 0.5%, Zr: 0% to 0.01% %, Cr: 4% to 12.5%, Mo: 0.4% to 4.5%, W: 5% to 10%, Re: 1% to 6%, Ta: 5% or more 12% or less, Nb: 0.3% to 3%, Ti: 0% to less than 0.4%, Al: 4.7% to 5.7%, Co: 0.5% to 5.0% A high-strength Ni-base superalloy characterized in that the following and the balance substantially consist of Ni.

By weight, C: 0.06% to 0.09%, B: 0.016% to 0.035%, Hf: 0.2% to 0.4%, Cr: 6.5% to 8 Mo: 0.4% or more and 1.0% or less, W: 5.5% or more and 9.5% or less, Re: 1.2% or more and 3.1% or less, Ta: 8% or more and 10% or less. %, Nb: 0.3% to 1.0%, Al: 4.7% to 5.4%, Co: 0.8% to 1.2%, Ti: 0% to 0.4% % And a balance substantially consisting of Ni.

By weight, C: 0.06% to 0.08%, B: 0.016% to 0.035%, Hf: 0.2% to 0.3%, Cr: 6.9% to 7 0.3% or less, Mo: 0.7% or more and 1.0% or less, W: 7.0% or more and 9.0% or less, Re: 1.2% or more and 1.6% or less, Ta: 8.5% 9.5% or less, Nb: 0.6% or more and 1.0% or less, Al: 4.9% or more and 5.2% or less, Co: 0.8% or more and 1.2% or less, and the balance is substantial. A high-strength Ni-base superalloy characterized by being made of Ni.

The high-strength Ni-base superalloy unidirectionally solidified casting according to any one of claims 1 to 4, which has been directionally solidified by a unidirectional solidification method.

The Ni-based superalloy according to claim 5,
The creep rupture life in the solidification direction should be 350 hours or more under conditions of 1040 ° C and 14 kgf / mm ² , and the creep rupture life in the direction perpendicular to the solidification direction should be 30 hours or more in the condition of 927 ° C and 32 kgf / mm ^2. Characterized high-strength Ni-base superalloy.

The Ni-based superalloy according to claim 5,
A solution heat treatment having a rectangular cubic γ 'phase with a side of 0.5 μm or less in a volume fraction of 50% or more and a creep rupture life in a direction perpendicular to the solidification direction of 927 ° C. and 32 kgf / mm ² A high-strength Ni-base superalloy characterized in that the tensile strength at 800 ° C. in a direction perpendicular to the solidification direction is 95 kgf / mm ² or more for 30 hours or more.

A high-strength Ni-base superalloy single crystal casting comprising the alloy composition according to claim 1.