JP3559681B2 - Steam turbine blade and method of manufacturing the same - Google Patents
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Description
【０００１】
【発明の属する技術分野】
本発明は、新規な蒸気タービン翼に係り、特に蒸気タービンブレードに関するものである。
【０００２】
【従来の技術】
従来、スチームタービンのブレード，ディスクには、１２Ｃｒ系のフェライト系耐熱鋼が使用されてきたが、スチームタービンの蒸気温度は効率向上のため、従来の６００℃未満の温度から近年は、６００〜６３０℃の温度に上昇しつつある。
【０００３】
このような蒸気温度の高温化に伴ない、一部オーステナイト系のγ′析出強化型超耐熱合金が使用されるようになってきた。
【０００４】
ところが、γ′析出強化型超耐熱合金は、フェライト系耐熱鋼より一段と高い高温強度を有するものの、熱膨張係数がフェライト系より高いため、他のフェライト系の部材との熱膨張差の問題、さらに熱疲労強度が劣る等の問題がある。そのため超耐熱合金の中では、フェライト系に近い比較的低い熱膨張係数を有するＭ２５２等の使用が検討されている。また低熱膨張超耐熱合金として、特開昭４７−１３３０２ 号，特開昭５３−６２２５号および特開昭５３−５８４２７ 号などが提案されている。
【０００５】
【発明が解決しようとする課題】
Ｍ２５２は、比較的低い熱膨張係数と、高い強度を有するが、一方高価なＣｏを約１０％も含むために非常に高価であるという問題がある。また、クリープ破断時の延性が比較的小さな値であるため、長時間使用後の切り欠き感受性が低下するおそれがある。
【０００６】
また特開昭４７−１３３０２ 号および特開昭５３−６２２５号に開示される合金は、低熱膨張合金として知られるいわゆるインバー合金と同じメカニズムで低熱膨張を得ている。すなわち、ＦｅとＮｉのバランスによりキュリー点を調整して、強磁性状態での低い熱膨張を利用している。
【０００７】
しかしながら、このタイプの合金の場合、Ｃｒ添加により熱膨張係数が増加するので、高温強度や耐酸化性を向上させる目的でＣｒを高めることができず、またＦｅ−Ｎｉ（またはＣｏ）のバランスが重要なために、相当量のＦｅを含有させる必要がある。したがって、本系統の合金の場合は、低Ｃｒ，高Ｆｅのため、高温強度や耐酸化性が劣り、耐熱用途に適さないという問題がある。さらに切り欠き感受性が高くクリープラプチャー試験において、ノッチ部で破断しやすいという欠点がある。
【０００８】
一方、特開昭５３−５８４２７ 号に開示される合金は、Ｍｏを多量に含むことにより、低い熱膨張係数が得られるが、Ｍｏに加えＮｂも含むことにより熱間加工性が低下する問題がある。
【０００９】
本発明は、フェライト系耐熱鋼に近い熱膨張係数を有しながら、フェライト系耐熱鋼を大幅に上回る高温強度と良好な耐酸化性、および切り欠き感受性ならびにラプチャー破断延性を兼備し、γ′析出強化型超耐熱合金からなる蒸気タービン翼とその製造方法を提供することである。
【００１０】
【課題を解決するための手段】
クリープ破断時の延性を良好な値とするために、ＡｌとＴｉは両方共に析出強化相であるγ′（ガンマプライム）相を形成する元素であるが、Ａｌの割合が高くなるほどクリープ破断時の延性が高くなることを見出し、強度とのバランスで最適割合としてＡｌ／（Ａｌ＋０．５６Ｔｉ）で表わされる値を０．４５〜０．８の範囲とした。また、熱処理による延性向上であり、本合金に溶体化処理後、８２０〜８８０℃の１段目時効処理を施すことにより、大幅に延性が向上することを見出した。
【００１１】
本発明は、重量％で、Ｃ：０．２％ 以下，Ｓｉ：１％以下，Ｍｎ：１％以下，Ｃｒ：１０〜２４％、およびＭｏ，Ｗの１種または２種をＭｏ＋（１／２）×Ｗ：５〜１７％，Ａｌ：０．５〜２％，Ｔｉ：１〜３．５％，Ｆｅ：１０％以下，Ｂ：０．０２％ 以下を含有し、Ｎｉが４８〜８０％であるＮｉ基超耐熱合金からなる蒸気タービン翼にある。
【００１２】
本発明は、重量％で、Ｃ：０．２％ 以下，Ｓｉ：１％以下，Ｍｎ：１％以下、Ｃｒ：１０〜２４％、およびＭｏ，Ｗの１種または２種をＭｏ＋（１／２）×Ｗ：５〜１７％，Ｃｏ：５％以下，Ｎｂ：１．０％以下，Ａｌ：０．５〜２％，Ｔｉ：１〜３．５％，Ｆｅ：１０％以下、およびＢ：０．０２％以下，Ｚｒ：０．２％ 以下の１種または２種を含有し、Ｎｉが４８〜８０％であるＮｉ基超耐熱合金からなる蒸気タービン翼にある。
【００１３】
さらに本発明の望ましい組成は、重量％で、Ｃ：０．０８％以下，Ｓｉ：０．５％以下，Ｍｎ：０．５％ 以下，Ｃｒ：１５〜２２％、およびＭｏ，Ｗの１種または２種をＭｏ＋（１／２）×Ｗ：５〜１２％，Ａｌ：１．０〜１．８％，Ｔｉ：１．２〜３．０％，Ｆｅ：２％以下、およびＢ：０．０２％以下，Ｚｒ：０．２％ 以下の１種または２種を含有し、さらにＡｌ／（Ａｌ＋０．５６Ｔｉ）で表わされる値が０．４５〜０．７０であり、Ｎｉが４８〜８０％を有するものである。本発明は、常温から６００℃までの平均熱膨張係数が、１３．８×１０ マイナス６乗／℃以下であり、かつ６００℃における引張強度が１０００Ｎ／ｍｍ２ 以上、および試験温度６５０℃、荷重応力：６８６Ｎ／ｍｍ２ の条件下で、切り欠き−平滑複合クリープラプチャー試験を行った後の破断寿命が５０時間以上で、かつ破断時の絞りが３０％以上であるＮｉ基合金からなる蒸気タービン翼にある。また、上記合金のうち、高強度と高延性とを同時に満足させるには、溶解後、熱間鍛造を行ったのち、９８０〜１０８０℃での溶体化処理を施し、次いで８２０〜８８０℃での第１段時効処理、および６００〜８００℃での第２段時効処理を行う製造方法を実施するのが好ましい。
【００１４】
本発明は動翼，静翼に用いることができるが、特に動翼に用いるのが好ましく、高圧タービン，中圧タービン及び高中圧一体タービンの初段に用いるのが好ましく、高圧タービンでは４段目まで用いることができる。高中圧タービンでは高圧部で４段目まで及び中圧部の初段と２段に用いられる。
【００１５】
以下に本発明合金の成分限定理由について述べる。
【００１６】
Ｃは、炭化物形成により結晶粒粗大化を防止する効果を有する。しかし、多すぎると、炭化物がストリンガー状に析出しやすくなり、加工方向に対する直角方向の延性が低下し、さらにＴｉと結合して炭化物を形成するため、本来Ｎｉと結び付いて析出強化相となるγ′を形成するＴｉ量が確保できなくなるため、Ｃは０．２％以下に限定する。望ましいＣの範囲は０．１５％以下であり、より望ましくは０．０８％以下であり、０．００５〜０．０５％がよい。
【００１７】
ＭｎとＳｉは、合金溶製時に脱酸剤として用いられるが、過度に含有すると熱間加工性の低下や使用時の靭性を損なうため、それぞれＭｎ：１％以下，Ｓｉ：１％以下に限定する。望ましくは、Ｍｎ，Ｓｉそれぞれ０．５％以下，より０．１％以下が好ましく、特に無添加が好ましく、０．０１％ 以下が最も好ましい。
【００１８】
Ｃｒは、基地に固溶して、合金の耐酸化性を向上させるとともに切り欠きラプチャー感受性を大幅に緩和させる効果を有する。１０％未満では、上記効果が得られず、また過度の添加は合金の塑性加工が困難となるため、Ｃｒは１０〜２４％に限定する。望ましいＣｒの範囲は１５〜２２％であり、最も１８〜２０％が好ましい。
【００１９】
ＭｏおよびＷは、合金の熱膨張係数を下げる効果があり、１種または２種を添加する。Ｍｏ＋（１／２）×Ｗ量で５％未満では、上記効果が得られず、また１７％を越えると、合金の塑性加工が困難となるため、ＭｏとＷの１種または２種をＭｏ＋（１／２）×Ｗで５〜１７％に限定する。ＭｏとＷの望ましい範囲はＭｏ＋（１／２）×Ｗで５〜１２％である。Ｍｏは５〜１７％、好ましくは５〜１２％、Ｗは３０％以下、好ましくは３〜１５％である。特に、Ｍｏ単独が好ましく、８〜１２％がよい。
【００２０】
Ａｌは、γ′相と呼ばれる金属間化合物（Ｎｉ３Ａｌ）を形成し、合金の高温強度を高めるために添加する。上記効果を得るため０．５％以上が必要であるが、２％を越えると熱間加工が困難となるのでＡｌは０．５〜２％ に限定する。望ましいＡｌの範囲は０．８〜１．８％より好ましくは０．８〜１．５％である。
【００２１】
Ｔｉは、Ａｌと共にγ′相（Ｎｉ３（Ａｌ，Ｔｉ））を形成する。Ａｌ単独のγ′相よりもＡｌ，Ｔｉからなるγ′の方が、さらに高い高温強度が得られる。
【００２２】
そのためＴｉは、１％以上が必要であるが、３．５％ を越えるとγ′相が不安定になり、また熱間加工性の面でも好ましくないので、１〜３％に限定する。望ましいＴｉの範囲は、１．２〜３．０％であり、より２．０〜３．０％が好ましい。前述のように、本合金においてＡｌとＴｉのバランスは重要である。γ′相中のＡｌの割合が多くなるほど、延性は向上するが、逆に強度は低下する。本発明合金においては、十分な延性を確保することが重要であり、γ′相中のＡｌの割合を原子量の比として表わすため、Ａｌ／（Ａｌ＋０．５６Ｔｉ）なる数値を設定した。この値が０．４５より低いと十分な延性が得られない。逆に０．８を越えると強度が不足する。更に、（Ｔｉ／Ａｌ）比は強度に影響を及ぼし、その比は
Ｆｅ量との関係によって前述のように定めることができる。
【００２３】
ＢおよびＺｒは、粒界を強化し、合金の高温における延性を高める効果があるため、１種または２種を添加する。しかし過度に添加すると、かえって熱間加工性を劣化させるため、Ｂは０．０２％以下，Ｚｒは０．２％以下に限定した。特に、Ｂは０．００２〜０．０１５％，Ｚｒ０．０１〜０．１％が好ましく、よりＢは
０．００４〜０．０１０％が好ましい。
【００２４】
Ｆｅは、必ずしも添加する必要はないが、合金の熱間加工性を改善する作用があるため、必要に応じて添加することができる。１０％を越えると、合金の熱膨張係数が大きくなり、また耐酸化性が劣化するため、上限を１０％に限定するのがよい。望ましくは２％以下であり、より無添加がよい。
【００２５】
Ｃｏは、合金に固溶して、合金の引張強度およびクリープ破断強度を向上させる効果があり、必要に応じて添加することができる。Ｃｏは高価な元素であるため添加する場合には、上限を５％とするのがよい。
【００２６】
Ｎｂは、ＡｌやＴｉとともにγ′相であるＮｉ３（Ａｌ，Ｔｉ，Ｎｂ）を形成し、高温強度向上に寄与するため、必要に応じて添加することができる。しかし、多すぎるとＮｉ２Ｎｂ を主体とするＬａｖｅｓ 相を形成しやすく、強度上昇に寄与しないばかりか延性も低下させる。特に多量のＭｏ、あるいはＭｏとＦｅ含有量が多い場合にはＬａｖｅｓ 相が形成されやすくなる。少量のＬａｖｅｓ 相の場合、熱処理等で消失させることも可能であるが、製造工程が繁雑となり好ましくない。したがって、Ｎｂを添加する場合でも、Ｎｂの上限は１．０％ が好ましい。より好ましいＮｂの上限は０．８％ であり、０．１〜０．７％が特によい。
【００２７】
Ｎｉ量は４８％未満ではＮｉ基合金が得られなくなり強度等が低下するので、４８％以上とする。また、８０％を越えると延性が低下するので、８０％以下とする。特に４９〜７５％が好ましく、より５４〜７０％が好ましい。
【００２８】
更に、Ｐ：０．０５％以下，Ｓ：０．０１％以下，Ｃｕ：５％以下，Ｍｇ：０．０１％以下，Ｃａ：０．０１％以下とするのが好ましい。
【００２９】
次に熱処理方法について述べる。本発明による合金は、熱処理条件によって、炭化物を粒界に析出させ、クリープ破断時の延性を向上させることができる。本発明者らは、本発明合金の熱処理条件について鋭意検討を行った結果、溶体化処理後、２段時効処理を行うことにより、炭化物を析出させて安定化させ、高温強度を劣化させることなく、安定した延性が得られる知見を得たものである。
【００３０】
上記組成範囲内の合金元素を適正に組み合わせることにより、常温から６００℃までの平均熱膨張係数が、１３．８×１０ マイナス６乗／℃以下の低熱膨張と、６００℃における引張強度が１０００Ｎ／ｍｍ２ 以上および試験温度６５０℃，荷重応力：６８６Ｎ／ｍｍ２ 条件下で、切り欠き−平滑複合クリープラプチャー試験を行った後の破断寿命が５０時間以上でかつ破断時の絞りが３０％以上である高い高温強度を兼備させることができる。
【００３１】
【発明の実施の形態】
〔実施例１〕
表１に示す組成の合金を、それぞれ１０ｋｇ真空溶解して造塊し、続いて３０ｍｍ角に熱間鍛造した。次いで２種類の熱処理を施した。熱処理Ａは、１０６６℃で４時間加熱後空冷し、さらに７２０℃で８時間加熱後、１時間に約５５℃の速度で６２０℃まで冷却し、さらに６２０℃で８時間加熱後空冷の熱処理である。次に熱処理Ｂは、第１段時効処理として、８５０℃で４時間加熱後、空冷し、第２段時効処理として、７６０℃で１６時間加熱後、空冷の熱処理である。なお、鍛造時に割れ等は発生せず、鍛造性は良好であった。さらに本発明合金と以下に示す特性を比較するため、従来合金（Ｍ２５２相当）も作製した。
【００３２】
表２に本発明合金，従来合金の常温から各温度までの平均熱膨張係数を示す。本発明合金が通常使用される温度は、６００〜７００℃であるが、２０℃から
６００℃および７００℃までの熱膨張係数は、従来合金とほぼ同様の、フェライト系耐熱鋼並みの低い熱膨張係数を示している。
【００３３】
【表１】
【表２】
表３に本発明合金，従来合金の常温における引張試験結果を、表４に６００℃における引張試験結果を示す。本発明合金は従来合金とほぼ同等の高い強度を示している。
【００３６】
表５に本発明合金のうちの１７合金を選び、試験温度：６５０℃，荷重応力：６８６Ｎ／ｍｍ２ の条件で切り欠き−平滑複合クリープラプチャー試験を行った結果を示す。表５から本発明合金は、すべて平滑部で破断し、切り欠き感受性も良好であり、また寿命も十分長いことがわかる。
【００３７】
また、本発明合金のうち、特にＡｌ／（Ａｌ＋０．５６Ｔｉ）の値が０．４５を越えるＮｏ．１４〜３１の延性が高い。さらに熱処理Ｂを行うことで、クリープ破断延性は、一段と向上しており、いずれの合金も３０％以上の絞りが出ているのがわかる。
【００３８】
しかし、Ａｌ／（Ａｌ＋０．５６Ｔｉ）値が０．７を越えるＮｏ．３１合金は、延性は高いものの、破断時間がやや低下している。したがって、良好なクリープ破断特性と強度を両立させるため、Ａｌ／（Ａｌ＋０．５６Ｔｉ）値を０．４５〜
０．７ に制限し、かつ熱処理Ｂを施すことが有効であることがわかる。
【００３９】
【表３】
【表４】
【表５】
〔実施例２〕
実施例１と同様に真空溶解によって表６（重量％）に示すＮｉ基合金の鋳塊を製造し、熱間鍛造によって所望の寸法の角材を得た。残部はＮｉである。
【００４３】
【表６】
図１は植込型式が鞍型タイプ，シュラウドチカバーがテノンシュラウドカバータイプである６２５℃級蒸気タービンの高圧タービンの初段動翼の斜視図である。図中、１は翼部、２はテノン、３はダブティルである。テノンは弯曲していて個々の動翼に対して個々のテノンが機械的にかしめによって結合される。
【００４５】
図２は植込型式がアキシャルタイプ，シュラウドカバーがテノン・ダブル・シュラウドタイプである同じ６２５℃級蒸気タービンの高圧タービンの初段動翼の斜視図である。テノンが複数の動翼を１個のテノンで機械的にかしめによって固定されるものである。
【００４６】
図３は同じく６２５℃級蒸気タービンの中圧タービンの初段動翼の斜視図である。植込型式は逆クリスマスツリータイプ，シュラウドカバーはテノン・シュラウドタイプであり、個々の動翼に個々のテノンがかしめによって固定したものである。
【００４７】
図４は同じく６２５℃級中圧タービンの初段動翼の斜視図であり、図３と植込型式及びシュラウドカバーが同じ形式のものである。本図面でのテノンは同じ素材から一体に形成されたものである。
【００４８】
本実施例は図１〜図４の動翼を表６に示す合金によって製造したものであり、前述の鍛造後に各図の形式に対して所望の相似形になるように熱間型鍛造を行った。次いで、実施例１の熱処理Ａ及びＢと同じ熱処理を施し、図１〜図４の各々の形状に機械加工した。図１〜図３のテノンは同じ組成のＮｉ基合金が好ましいが、高強度１２％Ｃｒ系マルテンサイト鋼でもよい。いずれも鍛造材である。
【００４９】
表７は室温（２０℃），６００℃での引張試験結果を示すものである。室温ではいずれも目標の耐力６９０Ｎ／ｍｍ２ 以上及び引張強さ９６０Ｎ／ｍｍ２ 以上を有しており、伸び率２０％以上，絞り率２５％以上の高いものである。また、６００℃では目標の耐力が３３５Ｎ／ｍｍ２ 以上，引張強さ７７０Ｎ／ｍｍ２ 以上有し、伸び率１０％以上，絞り率１５％以上の高い値を有している。
【００５０】
【表７】
表８は熱膨張係数（×１０−６／℃）を示し、室温から６００℃までの平均熱膨張係数１３．８×１０−６／℃以下である。
【００５２】
【表８】
表９は６５０℃における複合試験片を用いてクリープ破断試験したものである。表８に示すように、高応力下ではいずれも高い寿命を有しています。
【００５４】
【表９】
【発明の効果】
以上のように本発明の蒸気タービン翼は、室温から７００℃までの温度変化に対して、熱膨張係数が小さく、また６００℃における引張特性も良好で、かつ６５０℃におけるラプチャー寿命も十分長く、また破断時の延性も良好であるため従来のフェライト系の耐熱鋼より高い高温強度を有し、かつフェライト系に近い熱膨張係数を有したもので、蒸気温度として６００〜６５０℃の蒸気タービンの高温化に対応でき、高い熱効率を上げることができ、その効果は非常に大きい。
【図面の簡単な説明】
【図１】高圧タービンの初段動翼の斜視図。
【図２】高圧タービンの初段動翼の斜視図。
【図３】中圧タービンの初段動翼の斜視図。
【図４】中圧タービンの初段動翼の斜視図。
【符号の説明】
１…翼部、２…テノン、３…ダブティル。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a novel steam turbine blade, and particularly to a steam turbine blade.
[0002]
[Prior art]
Conventionally, 12Cr ferritic heat-resistant steel has been used for steam turbine blades and disks. However, the steam temperature of the steam turbine has been increased from the conventional temperature of less than 600 ° C. to 600 to 630 in recent years in order to improve efficiency. The temperature is rising to ° C.
[0003]
With the increase in the steam temperature, austenitic γ ′ precipitation-strengthened super heat-resistant alloys have been used in part.
[0004]
However, γ 'precipitation-strengthened superalloys have higher high-temperature strength than ferritic heat-resistant steel, but have a higher coefficient of thermal expansion than ferrite-based alloys. There are problems such as poor thermal fatigue strength. Therefore, among super heat-resistant alloys, use of M252 or the like having a relatively low coefficient of thermal expansion close to that of ferrite is under study. Japanese Patent Application Laid-Open Nos. 47-13302, 53-6225, and 53-58427 have been proposed as low thermal expansion super heat resistant alloys.
[0005]
[Problems to be solved by the invention]
M252 has a relatively low coefficient of thermal expansion and high strength, but has a problem that it is very expensive because it contains about 10% of expensive Co. In addition, since the ductility at the time of creep rupture is a relatively small value, the notch sensitivity after long-term use may be reduced.
[0006]
The alloys disclosed in JP-A-47-13302 and JP-A-53-6225 obtain low thermal expansion by the same mechanism as a so-called invar alloy known as a low thermal expansion alloy. That is, the Curie point is adjusted by the balance between Fe and Ni, and low thermal expansion in a ferromagnetic state is used.
[0007]
However, in the case of this type of alloy, the addition of Cr increases the coefficient of thermal expansion, so that Cr cannot be increased for the purpose of improving high-temperature strength and oxidation resistance, and the balance of Fe—Ni (or Co) is not improved. Because of its importance, it is necessary to contain a considerable amount of Fe. Therefore, in the case of the alloy of the present system, there is a problem that high-temperature strength and oxidation resistance are inferior due to low Cr and high Fe, and the alloy is not suitable for heat-resistant applications. Furthermore, there is a disadvantage that the notch portion is easily broken in a creep rupture test due to high notch sensitivity.
[0008]
On the other hand, the alloy disclosed in Japanese Patent Application Laid-Open No. 53-58427 can obtain a low coefficient of thermal expansion by containing a large amount of Mo, but has a problem that the hot workability is reduced by containing Nb in addition to Mo. is there.
[0009]
The present invention, while having a thermal expansion coefficient close to that of ferritic heat-resistant steel, combines high-temperature strength that is significantly higher than ferritic heat-resistant steel, good oxidation resistance, and notch sensitivity and rupture rupture ductility. An object of the present invention is to provide a steam turbine blade made of a reinforced super heat-resistant alloy and a method of manufacturing the same.
[0010]
[Means for Solving the Problems]
In order to obtain a good value of the ductility at the time of creep rupture, both Al and Ti are elements forming a γ ′ (gamma prime) phase which is a precipitation strengthening phase. It was found that the ductility was increased, and the value represented by Al / (Al + 0.56Ti) was set in the range of 0.45 to 0.8 as the optimum ratio in balance with the strength. It was also found that the ductility was improved by heat treatment, and that the ductility was significantly improved by subjecting the alloy to a first-stage aging treatment at 820 to 880 ° C. after the solution treatment.
[0011]
In the present invention, C: 0.2% or less, Si: 1% or less, Mn: 1% or less, Cr: 10 to 24%, and one or two of Mo and W are Mo + (1 / 2) × W: 5 to 17%, Al: 0.5 to 2%, Ti: 1 to 3.5%, Fe: 10% or less, B: 0.02% or less, Ni: 48 to 80 % Of a steam turbine blade made of a Ni-base super heat-resistant alloy.
[0012]
In the present invention, C: 0.2% or less, Si: 1% or less, Mn: 1% or less, Cr: 10 to 24%, and one or two of Mo and W are Mo + (1 / 2) × W: 5 to 17%, Co: 5% or less, Nb: 1.0% or less, Al: 0.5 to 2%, Ti: 1 to 3.5%, Fe: 10% or less, and B : 0.02% or less, Zr: 0.2% or less. A steam turbine blade comprising a Ni-based super heat-resistant alloy containing 48 to 80% of Ni.
[0013]
Further, a desirable composition of the present invention is, by weight%, C: 0.08% or less, Si: 0.5% or less, Mn: 0.5% or less, Cr: 15 to 22%, and one of Mo and W. Alternatively, Mo + (1/2) × W: 5 to 12%, Al: 1.0 to 1.8%, Ti: 1.2 to 3.0%, Fe: 2% or less, and B: 0 0.02% or less, Zr: 0.2% or less, and a value represented by Al / (Al + 0.56Ti) is 0.45 to 0.70, and Ni is 48 to 80. %. The present invention has an average thermal expansion coefficient from room temperature to 600 ° C. of 13.8 × 10 −6 / ° C. or less, a tensile strength at 600 ° C. of 1000 N / mm 2 or more, a test temperature of 650 ° C., and a load. A steam turbine made of a Ni-based alloy having a rupture life of 50 hours or more after performing a notch-smooth composite creep rupture test under a condition of stress: 686 N / mm 2 and a squeezing at break of 30% or more. On the wings. In order to simultaneously satisfy the high strength and the high ductility of the above alloys, after melting, hot forging is performed, and then a solution treatment is performed at 980 to 1080 ° C., and then at 820 to 880 ° C. It is preferable to carry out a manufacturing method in which a first-stage aging treatment and a second-stage aging treatment at 600 to 800 ° C. are performed.
[0014]
The present invention can be used for a moving blade and a stationary blade, but is particularly preferably used for a moving blade, and is preferably used for the first stage of a high-pressure turbine, a medium-pressure turbine, and a high-medium-pressure integrated turbine. Can be used. In a high-to-medium pressure turbine, it is used in the high pressure section up to the fourth stage and in the first and second stages of the medium pressure section.
[0015]
The reasons for limiting the components of the alloy of the present invention are described below.
[0016]
C has an effect of preventing crystal grain coarsening due to carbide formation. However, if it is too large, carbides tend to precipitate in a stringer-like manner, ductility in the direction perpendicular to the processing direction is reduced, and furthermore, since carbides are formed by combining with Ti, γ which is originally combined with Ni to form a precipitation strengthening phase C is limited to 0.2% or less because the amount of Ti forming '′ cannot be secured. The desirable range of C is 0.15% or less, more preferably 0.08% or less, and 0.005 to 0.05% is good.
[0017]
Mn and Si are used as deoxidizers during the melting of alloys. However, Mn and Si are limited to Mn: 1% or less and Si: 1% or less, respectively. I do. Desirably, Mn and Si are each preferably 0.5% or less, more preferably 0.1% or less, particularly preferably no addition, and most preferably 0.01% or less.
[0018]
Cr forms a solid solution in the matrix and has the effect of improving the oxidation resistance of the alloy and greatly reducing the notch rupture sensitivity. If the content is less than 10%, the above effects cannot be obtained, and if excessive addition makes the plastic working of the alloy difficult, the content of Cr is limited to 10 to 24%. The desirable range of Cr is 15 to 22%, most preferably 18 to 20%.
[0019]
Mo and W have the effect of lowering the coefficient of thermal expansion of the alloy, and one or two of them are added. If the amount of Mo + (1/2) × W is less than 5%, the above effect cannot be obtained. If the amount exceeds 17%, plastic working of the alloy becomes difficult. Therefore, one or two types of Mo and W are used as Mo +. (1/2) × W is limited to 5 to 17%. A desirable range of Mo and W is Mo + (1/2) × W, which is 5 to 12%. Mo is 5 to 17%, preferably 5 to 12%, and W is 30% or less, preferably 3 to 15%. Particularly, Mo alone is preferable, and 8 to 12% is preferable.
[0020]
Al forms an intermetallic compound (Ni 3 Al) called a γ ′ phase and is added to increase the high-temperature strength of the alloy. To obtain the above effects, 0.5% or more is necessary. However, if it exceeds 2%, hot working becomes difficult, so the Al content is limited to 0.5 to 2%. A desirable range of Al is 0.8 to 1.8%, more preferably 0.8 to 1.5%.
[0021]
Ti forms a γ ′ phase (Ni 3 (Al, Ti)) together with Al. Γ ′ composed of Al and Ti can obtain higher high-temperature strength than γ ′ phase of Al alone.
[0022]
Therefore, 1% or more of Ti is required. However, if it exceeds 3.5%, the γ 'phase becomes unstable, and it is not preferable in terms of hot workability. Desirable range of Ti is 1.2 to 3.0%, more preferably 2.0 to 3.0%. As described above, in the present alloy, the balance between Al and Ti is important. As the proportion of Al in the γ 'phase increases, ductility improves, but conversely, strength decreases. In the alloy of the present invention, it is important to ensure sufficient ductility, and a numerical value of Al / (Al + 0.56Ti) is set in order to express the ratio of Al in the γ 'phase as a ratio of atomic weight. If this value is lower than 0.45, sufficient ductility cannot be obtained. Conversely, if it exceeds 0.8, the strength will be insufficient. Further, the (Ti / Al) ratio affects the strength, and the ratio can be determined as described above in relation to the amount of Fe.
[0023]
Since B and Zr have the effect of strengthening the grain boundaries and increasing the ductility of the alloy at high temperatures, one or two of them are added. However, if added excessively, the hot workability is rather deteriorated, so that B is limited to 0.02% or less and Zr is limited to 0.2% or less. In particular, B is preferably 0.002 to 0.015% and Zr 0.01 to 0.1%, and more preferably B is 0.004 to 0.010%.
[0024]
Fe does not always need to be added, but can improve the hot workability of the alloy, and can be added as needed. If it exceeds 10%, the coefficient of thermal expansion of the alloy increases and the oxidation resistance deteriorates. Therefore, it is preferable to limit the upper limit to 10%. Desirably, it is 2% or less, and it is better not to add it.
[0025]
Co is dissolved in the alloy and has an effect of improving the tensile strength and creep rupture strength of the alloy, and can be added as necessary. Since Co is an expensive element, the upper limit is preferably set to 5% when Co is added.
[0026]
Nb forms Ni 3 (Al, Ti, Nb), which is a γ ′ phase, together with Al and Ti, and contributes to improvement in high-temperature strength. Therefore, Nb can be added as necessary. However, when the content is too large, a Laves phase mainly composed of Ni 2 Nb is easily formed, and not only does not contribute to an increase in strength but also decreases ductility. In particular, when a large amount of Mo or a large content of Mo and Fe is used, a Laves phase is easily formed. In the case of a small amount of Laves phase, it can be eliminated by heat treatment or the like, but this is not preferable because the production process becomes complicated. Therefore, even when Nb is added, the upper limit of Nb is preferably 1.0%. The more preferable upper limit of Nb is 0.8%, and 0.1 to 0.7% is particularly preferable.
[0027]
If the Ni content is less than 48%, a Ni-based alloy cannot be obtained, and strength and the like are reduced. Therefore, the Ni content is set to 48% or more. On the other hand, if it exceeds 80%, the ductility is reduced. Particularly, it is preferably from 49 to 75%, more preferably from 54 to 70%.
[0028]
Further, it is preferable that P: 0.05% or less, S: 0.01% or less, Cu: 5% or less, Mg: 0.01% or less, and Ca: 0.01% or less.
[0029]
Next, a heat treatment method will be described. The alloy according to the present invention can improve the ductility at the time of creep rupture by precipitating carbides at the grain boundaries depending on the heat treatment conditions. The present inventors have conducted intensive studies on the heat treatment conditions of the alloy of the present invention. As a result of performing a two-stage aging treatment after the solution treatment, the carbides are precipitated and stabilized, without deteriorating the high-temperature strength. It has been found that stable ductility can be obtained.
[0030]
By appropriately combining alloy elements within the above composition range, the average thermal expansion coefficient from room temperature to 600 ° C. is low thermal expansion of 13.8 × 10 −6 / ° C. or less, and the tensile strength at 600 ° C. is 1000 N / The notch-smooth composite creep rupture test was performed for a fracture life of 50 hours or more and a reduction of 30% or more at the time of fracture under a condition of not less than 2 mm, a test temperature of 650 ° C., and a load stress of 686 N / mm 2. Some high high-temperature strength can be provided.
[0031]
BEST MODE FOR CARRYING OUT THE INVENTION
[Example 1]
Each of the alloys having the compositions shown in Table 1 was melted in a vacuum of 10 kg to form an ingot, followed by hot forging into a 30 mm square. Next, two types of heat treatment were performed. Heat treatment A is heated at 1066 ° C. for 4 hours, air-cooled, further heated at 720 ° C. for 8 hours, cooled to 620 ° C. at a rate of about 55 ° C. per hour, further heated at 620 ° C. for 8 hours, and then air-cooled. is there. Next, the heat treatment B is a first-stage aging treatment of heating at 850 ° C. for 4 hours, followed by air cooling, and a second-stage aging treatment of 760 ° C. for 16 hours, followed by air cooling. In addition, cracking did not occur at the time of forging, and the forgeability was good. Further, in order to compare the characteristics shown below with the alloy of the present invention, a conventional alloy (M252 equivalent) was also prepared.
[0032]
Table 2 shows the average thermal expansion coefficients of the alloy of the present invention and the conventional alloy from room temperature to each temperature. The temperature at which the alloy of the present invention is usually used is from 600 to 700 ° C., but the coefficient of thermal expansion from 20 ° C. to 600 ° C. and 700 ° C. is as low as that of the conventional alloy, similar to ferritic heat-resistant steel. The coefficient is shown.
[0033]
[Table 1]
[Table 2]
Table 3 shows the results of the tensile test at room temperature of the alloy of the present invention and the conventional alloy, and Table 4 shows the results of the tensile test at 600 ° C. The alloy of the present invention has almost the same high strength as the conventional alloy.
[0036]
Table 5 shows the results of a notch-smooth composite creep rupture test performed on 17 alloys of the present invention under the conditions of a test temperature: 650 ° C. and a load stress: 686 N / mm 2 . From Table 5, it can be seen that all the alloys of the present invention break at the smooth portion, have good notch sensitivity, and have a sufficiently long life.
[0037]
Further, among the alloys of the present invention, in particular, the alloys of No. 1 in which the value of Al / (Al + 0.56Ti) exceeds 0.45. The ductility of 14 to 31 is high. Further, by performing the heat treatment B, the creep rupture ductility is further improved, and it can be seen that the reduction of 30% or more has been achieved in all the alloys.
[0038]
However, when the Al / (Al + 0.56Ti) value exceeds 0.7, the Al. The 31 alloy has a high ductility, but the rupture time is slightly reduced. Therefore, in order to achieve both good creep rupture characteristics and strength, the Al / (Al + 0.56Ti) value is set to 0.45 to 0.45.
It is understood that it is effective to limit the value to 0.7 and to perform the heat treatment B.
[0039]
[Table 3]
[Table 4]
[Table 5]
[Example 2]
Ingots of the Ni-based alloys shown in Table 6 (% by weight) were produced by vacuum melting in the same manner as in Example 1, and square bars having desired dimensions were obtained by hot forging. The balance is Ni.
[0043]
[Table 6]
FIG. 1 is a perspective view of a first stage rotor blade of a high-pressure turbine of a 625 ° C. class steam turbine in which an implantation type is a saddle type and a shroud cover is a Tenon shroud cover type. In the figure, 1 is a wing, 2 is a tenon, and 3 is a dovetil. The tenon is curved and the individual tenons are mechanically coupled to the individual blades by caulking.
[0045]
FIG. 2 is a perspective view of a first stage rotor blade of a high-pressure turbine of the same 625 ° C. class steam turbine in which an implantation type is an axial type and a shroud cover is a Tenon double shroud type. The tenon mechanically fixes a plurality of rotor blades with one tenon.
[0046]
FIG. 3 is a perspective view of the first stage rotor blade of the 625 ° C. class steam turbine medium pressure turbine. The implantation type is an inverted Christmas tree type, and the shroud cover is a tenon shroud type. Each tenon is fixed to each bucket by caulking.
[0047]
FIG. 4 is a perspective view of a first stage rotor blade of the 625 ° C. class medium pressure turbine, in which the implantation type and the shroud cover are the same as those in FIG. 3. Tenon in this drawing is formed integrally from the same material.
[0048]
In this embodiment, the rotor blades shown in FIGS. 1 to 4 are manufactured by using the alloys shown in Table 6, and after the above-described forging, hot die forging is performed so as to have a desired similar shape to the type shown in each drawing. Was. Next, the same heat treatment as that of heat treatments A and B of Example 1 was performed, and machining was performed to each of the shapes shown in FIGS. 1 to 3 are preferably Ni-based alloys having the same composition, but may be high-strength 12% Cr-based martensitic steel. All are forged materials.
[0049]
Table 7 shows the results of the tensile test at room temperature (20 ° C.) and 600 ° C. At room temperature, all have the target proof stress of 690 N / mm 2 or more and the tensile strength of 960 N / mm 2 or more, and have high elongation of 20% or more and draw ratio of 25% or more. At 600 ° C., the target proof stress is 335 N / mm 2 or more, the tensile strength is 770 N / mm 2 or more, and the elongation is 10% or more and the draw ratio is 15% or more.
[0050]
[Table 7]
Table 8 shows the coefficient of thermal expansion (× 10 −6 / ° C.), which is not more than 13.8 × 10 −6 / ° C. from room temperature to 600 ° C.
[0052]
[Table 8]
Table 9 shows the results of a creep rupture test using a composite test piece at 650 ° C. As shown in Table 8, all have high life under high stress.
[0054]
[Table 9]
【The invention's effect】
As described above, the steam turbine blade of the present invention has a small coefficient of thermal expansion, a good tensile property at 600 ° C., and a sufficiently long rupture life at 650 ° C. with respect to a temperature change from room temperature to 700 ° C. In addition, since the ductility at break is also good, it has a high temperature strength higher than that of conventional ferritic heat-resistant steel, and has a thermal expansion coefficient close to that of ferritic steel. It can cope with high temperature, can raise high thermal efficiency, and its effect is very large.
[Brief description of the drawings]
FIG. 1 is a perspective view of a first-stage bucket of a high-pressure turbine.
FIG. 2 is a perspective view of a first-stage bucket of a high-pressure turbine.
FIG. 3 is a perspective view of a first stage rotor blade of the intermediate pressure turbine.
FIG. 4 is a perspective view of a first stage rotor blade of the intermediate pressure turbine.
[Explanation of symbols]
1 ... wings, 2 ... tenon, 3 ... dovel.
Claims (5)
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP12372597A JP3559681B2 (en) | 1997-05-14 | 1997-05-14 | Steam turbine blade and method of manufacturing the same |
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| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP12372597A JP3559681B2 (en) | 1997-05-14 | 1997-05-14 | Steam turbine blade and method of manufacturing the same |
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| Publication Number | Publication Date |
|---|---|
| JPH10317079A JPH10317079A (en) | 1998-12-02 |
| JP3559681B2 true JP3559681B2 (en) | 2004-09-02 |
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| JP12372597A Expired - Lifetime JP3559681B2 (en) | 1997-05-14 | 1997-05-14 | Steam turbine blade and method of manufacturing the same |
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| WO2010038680A1 (en) | 2008-09-30 | 2010-04-08 | 日立金属株式会社 | Process for manufacturing ni-base alloy and ni-base alloy |
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| US7160400B2 (en) * | 1999-03-03 | 2007-01-09 | Daido Tokushuko Kabushiki Kaisha | Low thermal expansion Ni-base superalloy |
| DE10035224A1 (en) * | 2000-07-20 | 2002-01-31 | Man Turbomasch Ag Ghh Borsig | Process and blank for producing rhomboid blades for axial turbomachines |
| JP4509664B2 (en) * | 2003-07-30 | 2010-07-21 | 株式会社東芝 | Steam turbine power generation equipment |
| ITMI20042482A1 (en) * | 2004-12-23 | 2005-03-23 | Nuovo Pignone Spa | STEAM TURBINE |
| ITMI20042483A1 (en) * | 2004-12-23 | 2005-03-23 | Nuovo Pignone Spa | STEAM TURBINE |
| JP4543380B2 (en) * | 2004-12-24 | 2010-09-15 | 日立金属株式会社 | Fuel cell stack fastening bolt alloy |
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| JP4972972B2 (en) * | 2006-03-22 | 2012-07-11 | 大同特殊鋼株式会社 | Ni-based alloy |
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| JP4768672B2 (en) * | 2007-05-30 | 2011-09-07 | 株式会社日本製鋼所 | Ni-base alloy excellent in structure stability and high-temperature strength and method for producing Ni-base alloy material |
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| JP2010249050A (en) * | 2009-04-16 | 2010-11-04 | Toshiba Corp | Steam turbine and steam turbine installation |
| JP5657964B2 (en) * | 2009-09-15 | 2015-01-21 | 三菱日立パワーシステムズ株式会社 | High-strength Ni-base forged superalloy and manufacturing method thereof |
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-
1997
- 1997-05-14 JP JP12372597A patent/JP3559681B2/en not_active Expired - Lifetime
Cited By (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| EP2128283A2 (en) | 2008-05-21 | 2009-12-02 | Kabushiki Kaisha Toshiba | Nickel-base casting superalloy and cast component for steam turbine using the same |
| US9238853B2 (en) | 2008-05-21 | 2016-01-19 | Kabushiki Kaisha Toshiba | Nickel-base casting superalloy and cast component for stream turbine using the same as material |
| WO2010038680A1 (en) | 2008-09-30 | 2010-04-08 | 日立金属株式会社 | Process for manufacturing ni-base alloy and ni-base alloy |
| US8845958B2 (en) | 2008-09-30 | 2014-09-30 | Hitachi Metals, Ltd. | Process for manufacturing Ni-base alloy and Ni-base alloy |
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| Publication number | Publication date |
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| JPH10317079A (en) | 1998-12-02 |
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