JP4168643B2 - Compound for bonded magnet and method for manufacturing the same - Google Patents

Description

【０２０５】
本発明によるナノコンポジット磁粉は、その組成および組織の特徴のため、その磁気特性に粒径依存性が小さいという特徴を有している。ナノコンポジット磁粉は、希土類元素Ｒの含有率が比較的低く、Ｒがリッチな粒界相が存在しないのに加え、Ｒ₂Ｆｅ₁₄Ｂ相を取り囲むように小さな硼化物相が分散しており、さらにＴｉは硼素との親和性が高いので硼化物相は他の相よりも多くのＴｉを含有している。その結果、ナノコンポジット磁粉は、従来の急冷磁石粉末に比べ耐酸化性に優れており、優れた磁気特性を粉砕後に維持できる。
【０２０６】
従来の急冷磁石粉末は比較的多量の希土類元素Ｒを含むので酸化されやすく、粒径が小さいほど粉末粒子表面の酸化による磁気特性の低下が顕著となる。例えば、ＭＱＰ−Ｂ（最大粒径３００μｍ以下）では、表２に示すように、粒径が７５μｍ以下、特に５３μｍ以下の粉末粒子の磁気特性が低下している。残留磁束密度Ｂ_rについてみると、最も高い値を示している１２５μｍ超１５０μｍ以下の粉末粒子の残留磁束密度Ｂ_r（０．９０Ｔ）に対して、５３μｍ以下の粉末粒子の残留磁束密度Ｂ_r（０．７９Ｔ）は９０％未満にまで低下している。また、（ＢＨ）ｍａｘについて見ると、５３μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘ（３８μｍ以下と３８μｍ超５３μｍ以下の値の単純平均）は８５．５ｋＪ／ｍ³であり、１５０μｍ超２１２μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘ（１５０μｍ超１８０μｍ以下と１８０μｍ超２１２μｍ以下の値の単純平均）である１１４．６ｋＪ／ｍ³の７５％未満にまで低下している。
【０２０７】
これに対し、ナノコンポジット磁粉は酸化による磁気特性の低下の割合が低く、磁気特性の粒径依存性が小さい。例えば、ナノコンポジット磁粉ＮＣＰ−０（最大粒径３００μｍ以下）では、表３に示すように、磁気特性はほとんど粒径に依存せず、優れた磁気特性を有している。例えば、残留磁束密度Ｂ_rは、最も高い値を示している１０６μｍ超１２５μｍ以下の粉末粒子の残留磁束密度Ｂ_r（０．８４５Ｔ）に対して、５３μｍ以下の粉末粒子の残留磁束密度Ｂ_r（約０．８２９Ｔ）は９８％以上の値を有している。また、（ＢＨ）ｍａｘについも、５３μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘは１０４．６ｋＪ／ｍ³であり、１５０μｍ超２１２μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘである１０６．６ｋＪ／ｍ³の９８％以上の値を有している。種々の組成のナノコンポジット磁粉について同様の評価を行った結果、ほとんどの組成についてナノコンポジット磁粉の５３μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘは、１５０μｍ超２１２μｍ以下の粉末粒子の平均の（ＢＨ）ｍａｘの９０％以上の値を有し、多くの組成について９５％以上の値が得られることが分かった。なお、磁粉の粒度分布の評価は、ＪＩＳ８８０１準拠の標準ふるいを用いて行った。
【０２０８】
【表２】

【０２０９】
【表３】

【０２１０】
このように、ナノコンポジット磁粉は従来の急冷磁石粉末と同等以上の磁気特性を有しているので、従来の急冷磁石粉末（例えばＭＱ粉）の代わりにボンド磁石用磁粉として用いることができる。勿論、ボンド磁石用磁粉をナノコンポジット磁粉のみで構成しても良いが、例えば、上述したＭＱ粉のうちの粒径が５３μｍ以下の粉末粒子をナノコンポジット磁粉に置き換えてもよい。
【０２１１】
以下に、５３μｍ以下および３８μｍ以下の微粒子を混入することによって充填性が改善される効果を実験結果を例示しながら説明する。
【０２１２】
まず、表４に示すような種々の粒度分布を有するナノコンポジット磁粉の試料ＮＣＰ−１からＮＣＰ−５を作製した。なお、ＮＣＰ−１の磁粉は、０．５ｍｍφのスクリーンを用いてパワーミルで粉砕することによって調製し、他のＮＣＰ−２〜ＮＣＰ−５の磁粉は、上述したピンミル装置を用いて、それぞれ回転数を３０００ｒｐｍ、４０００ｒｐｍ、５０００ｒｐｍおよび８０００ｒｐｍとすることによって調製した。これらの磁粉試料ＮＣＰ−１からＮＣＰ−５をタップデンサを用いてタップ密度を測定した結果を表５に示す。表５には、それぞれの磁粉試料中に含まれる粒径が５３μｍ以下の粉末粒子の質量％および粒径が２５０μｍ超の粉末粒子の質量％を合せて示している。
【０２１３】
表５の結果からわかるように、粒径が５３μｍ以下の粒子を１０質量％以上（厳密には９．５質量％以上）含む試料ＮＣＰ−３〜ＮＣＰ−５は、タップ密度が４．３ｇ／ｃｍ³以上と高く、磁粉の充填性が優れていることが分かる。磁粉のタップ密度で評価される磁粉の充填性は、ボンド磁石用のコンパウンドの粉末の充填性と相関しており、充填性の高い磁粉を用いて調製されたコンパウンドの粉末の充填性も高くなる。従って、粒径が５３μｍ以下のナノコンポジット磁粉を１０質量％含む磁粉を用いることによって、ボンド磁石用コンパウンドの粉末の充填性や流動性が改善され、高品質の成形体を得ることができる。
【０２１４】
【表４】

【０２１５】
【表５】

【０２１６】
さらに、成形密度を向上するためには、粒径が３８μｍ以下の粉末粒子を含むことが好ましい。表６に示す粒度分布を有するナノコンポジット磁粉ＮＣＰ−１１からＮＣＰ−１６を調製し、それぞれ２質量％のエポキシ樹脂と混合することによってコンパウンドを得た。それぞれのコンパウンドを用いて成形圧力９８０ＭＰａ（１０ｔ／ｃｍ²）で圧縮成形することによってボンド磁石成形体を得た。それぞれのボンド磁石成形体の密度を、それぞれのコンパウンドに用いた磁粉中の粒径が３８μｍ以下の粉末粒子の含有率とともに図１６に示す。
【０２１７】
【表６】

【０２１８】
図１６からわかるように、３８μｍ以下の粉末粒子の含有率が低すぎても高すぎても成形体の密度は低下する。種々検討した結果、十分な成形体密度を得るためには、粒径が３８μｍ以下の粉末粒子を約８質量％以上含む磁粉を用いることが好ましい。但し、粒径が３８μｍ以下の粉末粒子の含有率が約１６質量％を超える磁粉を用いると、成形性が低下し、高い密度の高品位の成形体が得られないことがある。
【０２１９】
〔コンパウンドおよび磁石体の製造方法の説明〕
上述のナノコンポジット磁粉を含むボンド磁石用磁粉は、樹脂等の結合剤と混合され、ボンド磁石用コンパウンドが製造される。
【０２２０】
射出成形用のコンパウンドは、公知の混練装置（例えばニーダや押出し機）を用いて磁粉と熱可塑性樹脂とを混練することによって製造される。また、圧縮成形用のコンパウンドは、溶剤で希釈した熱硬化性樹脂と磁粉とを混合し、溶剤を除去することによって製造される。得られた磁粉と樹脂との混合物は、必要に応じて、所定の粒度となるように解砕される。解砕の条件などを調整することによって、顆粒状としてもよい。また、粉砕によって得られた粉末材料を造粒してもよい。
【０２２１】
磁粉の耐食性を向上するために、磁粉の表面に予め化成処理等の公知の表面処理を施しても良い。さらに、磁粉の耐食性や樹脂との濡れ性、コンパウンドの成形性をさらに改善するために、シラン系、チタネート系、アルミネート系、ジルコネート系などの各種カップリング剤、コロイダルシリカなどセラミックス超微粒子、ステアリン酸亜鉛やステアリン酸カルシウムなどの潤滑剤を使用してもよく、熱安定剤、難燃剤、可塑剤などを使用してもよい。
【０２２２】
磁石用コンパウンドは種々の成形方法で種々の用途に用いられるので、用途に応じて、樹脂の種類および磁粉の配合比率が適宜決められる。樹脂としては、例えばエポキシ樹脂、フェノール樹脂やメラミン樹脂などの熱硬化性樹脂や、ポリアミド（ナイロン６６、ナイロン６、ナイロン１２等）や、ポリエチレン、ポリプロピレン、ポリ塩化ビニル、ポリエステル、ポリフェニレンサルファイドなどの熱可塑性樹脂や、ゴムやエラストマ、さらには、これらの変性体、共重合体、混合物などを用いることができる。
【０２２３】
さらに、本発明の磁粉によってコンパウンドの充填性および／または成形性が改善されるので、従来は用いることが難しかった高粘度の樹脂を用いることもできる。さらに、磁粉は酸化されにくいので、融点または軟化点が高く従来は使用できなかった樹脂（例えば、ポリイミドや液晶ポリマなど、また、種々の樹脂の高分子量グレード品）を用いることができるので、ボンド磁石の特性（耐熱性など）を改善することが出来る。また、熱硬化性樹脂を用いる場合においても、従来よりも高い温度で硬化する樹脂を用いることができる。
【０２２４】
成形方法としては、圧縮成形、圧延成形、カレンダー成形、押出し成形および射出成形を例示することができる。これらの成形方法のうち、圧縮成形、圧延成形および押出し成形では、比較的単純な形状の成形体しか成形できないが、成形時にあまり高い流動性が要求されないので、磁石粉末の充填率を高くできる。本発明による磁粉を用いることによって、従来よりも更に高い（例えば８０％を超える）磁粉充填率を実現することができ、最大で９０％程度まで充填することができる。但し、充填率を上げすぎると磁粉同士を十分に結合するための樹脂が不足し、ボンド磁石の機械的な強度の低下や、使用時に磁粉の脱落が生じる恐れがあるので、磁粉充填率は、８５％以下が好ましい。また、圧縮成形においては本発明の磁粉を用いることによって、成形体の表面に形成される空隙（ボイド）の量を減少でき、表面に形成する樹脂被膜への悪影響を抑制できるという利点が得られる。これらの成形方法には、適宜、熱硬化性樹脂、熱可塑性樹脂、ゴム等が用いられる。
【０２２５】
本発明による磁粉を用いると流動性が向上するので、特に、射出成形用コンパウンドに好適に用いられる。従来の急冷磁石粉末を用いたコンパウンドでは成形が困難であった複雑な形状の成形体を得ることができる。また、従来よりも高い充填率（例えば６５％を超える）で磁石粉末を配合できるので、磁石体の磁気特性を向上することができる。さらに、本発明による磁粉は、希土類元素の含有率が比較的少ないので、酸化され難い。従って、比較的軟化点の高い熱可塑性樹脂や熱可塑性エラストマを用いて、比較的高い温度で射出成形を行っても磁気特性が低下しない。
【０２２６】
［ボンド磁石の応用例］
本発明によるボンド磁石用コンパウンドは、上述したように、従来の急冷磁石粉末（例えばＭＱＩ社製の製品名ＭＱＰ−Ｂ）を用いたコンパウンドに比べ、優れた充填性（成形性）を有するとともに、耐熱性に優れており、且つ、従来の急冷磁石粉末を用いたボンド磁石と同等以上の磁気特性を有するボンド磁石を形成することができるので、種々の用途に好適に用いられる。
【０２２７】
図１７を参照しながら、ステッピングモータに応用した例を説明する。
【０２２８】
図１７は、永久磁石回転子型を備えるステッピングモータ１００の構成を模式的に示す分解斜視図である。ステッピングモータ１００は、ロータ１０１と、ロータ１０１の周辺に設けられたステータ部１０２とを有している。ロータ１０１は、外径８ｍｍの外周面を１０極に均等に着磁したボンド磁石を備えている。ステータ部１０２は、外ヨーク１０２ａおよび１０２ｂと、これらと互いに背中合わせに接合された２個の内ヨーク１０３と、これらの間に収容された励磁コイル１０４ａおよび１０４ｂとを備えている。このステッピングモータ１００は、１パルス電流に対応する励磁コイル１０４ａおよび１０４ｂの起磁力により１ステップ角だけロータ１０１が変位する動作を行う、いわゆるＰＭ型パルスモータである。
【０２２９】
ロータ１０１が備えるボンド磁石は、上述した本発明による充填性（成形性）に優れたコンパウンドを用いて形成されており、従来の急冷磁石粉末を用いたボンド磁石と同等以上の磁気特性を有するとともに、機械的特性に優れ、欠けなどが発生する恐れがなく、信頼性に優れている。また、耐熱性にも優れる。
【０２３０】
このような本発明によるコンパウンドを用いて形成されたボンド磁石を備えるステッピングモータは、小型・高性能で且つ信頼性に優れており、プリンターやディスクドライブ装置などのＯＡ機器やカメラやビデオなどのＡＶ機器などに好適に用いられる。
【０２３１】
ロータ１０１は、種々の方法で製造することができる。例えば、熱硬化性樹脂を用いたコンパウンドを圧縮成形することによって形成しても良いし、熱可塑性樹脂を用いたコンパウンドを射出成形または押出し成形することによって形成してもよい。以下、図１８を参照しながら、ロータ１０１の製造方法を説明する。
【０２３２】
例えば、熱硬化性樹脂を結合剤とするコンパウンドを用いる場合、図１８（ａ）〜（ｃ）を参照しながら説明するような成形方法を採用することによって、図１８（ｄ）に示すボンド磁石一体成形型のロータ２００を作製することができる。
【０２３３】
図１８（ｄ）に示したロータ２００は、ロータ軸２０５と、そのまわりに設けられたヨーク２０８と、ボンド磁石２１０とを備えている。ボンド磁石２１０は、ヨーク２０８の表面２１２に接着されている。
【０２３４】
ロータ２００は、図１８（ａ）から（ｃ）に示した工程で製造される。
【０２３５】
まず、図１８（ａ）に示したように、まず、粉末状のコンパウンド２０１を収容したフィーダボックス２０３を金型２０４の上面で摺動させながらコンパウンド２０１をキャビティ２０２内に充填する。金型２０４には、ロータ軸２０５がその中央に圧入されたヨーク２０８がセットされており、ロータ軸２０５を覆うように補助部材２０７が設けられている。金型２０４とこれらの間にキャビティ２０２が形成されている。
【０２３６】
次に、例えば、図１８（ｂ）に示すように、上パンチ２０９を介して、コンパウンド２０１を圧縮成形することによって、ヨーク２０８とコンパウンド２０１の成形体とを物理的に結合させる。
【０２３７】
次に、図１８（ｃ）に示すように、ロータ成形体を金型２０４から取り出す。補助部材は２０７は、ロータ軸２０５およびヨーク２０８から簡単に取り外され、ロータ軸２０５、ヨーク２０８、ボンド磁石２１０は一体化されている。但し、この状態では、ボンド磁石２１０はコンパウンドの粉末成形体であり、コンパウンドに含まれている熱硬化性樹脂は未硬化である。
【０２３８】
次に、ボンド磁石２１０を硬化するため、およびヨーク２０８とボンド磁石２１０との界面２１２における接合を強化するために、コンパウンドを所定の温度で硬化させる。硬化温度および硬化時間は用いる樹脂に応じて適宜設定される。
【０２３９】
本発明によるコンパウンドは、耐熱性に優れる磁粉を含んでいるので、従来よりも高い硬化温度で好適に硬化されるコンパウンドであり得る。従って、従来よりも、耐熱性、機械特性および接着強度に優れたボンド磁石２１０を形成することができる。さらに、本発明によるコンパウンドは、磁粉自体が耐食性に優れているため、熱硬化処理を大気中で行っても磁石特性の劣化は極めて小さい。従って、熱硬化処理を不活性雰囲気で行う必要がないので工程費用を削減できる。
【０２４０】
上述した成形方法によると、リング状のボンド磁石２１０を成形しながら、同時に、ヨーク２０８およびロータ軸２０５とボンド磁石２１０とを一体に成形できるので、ロ−タ２００を高い生産性で製造することができる。
【０２４１】
なお、金型２０４から成形体を取り出してから硬化する例を説明したが、金型２０４に加熱機構を設けて、金型２０４内で硬化してもよく、加圧した状態で硬化してもよい。さらに、圧縮成形に限られず、射出成形によってボンド磁石一体成形型ロータを形成することもできる。
【０２４２】
また、本発明によるコンパウンドは、従来の急冷磁石粉末を用いたコンパウンドに比べて高い充填性（成形性および／または流動性）を有するので、小さな間隙（例えば、約２ｍｍ幅）に確実に充填することができる。従って、本発明によるコンパウンドは、ＩＰＭ（Ｉｎｔｅｒｉｏｒ Ｐｅｒｍａｎｅｎｔ Ｍａｇｎｅｔ）型モータに用いられる磁石埋設型ロータ３００（図１９参照）の製造に好適に用いられる。
【０２４３】
図１９に示した磁石埋設型ロータ３００は、鉄心（例えば直径８０ｍｍ、厚さ５０ｍｍ）３０１と、鉄心３０１の中心に形成された回転軸スロット３０２と、鉄心３０１の周辺部に形成された複数のアーク状磁石スロット３０４とを備えている。ここでは、８個のアーク状磁石スロット３０４が設けられており、それぞれのスロット３０４は、第１スロット（例えば幅３．５ｍｍ）３０４ａと第２スロット（例えば幅１．８ｍｍ）３０４ｂとを有する２層構造となっている。これらのスロット３０４ａおよび３０４ｂ内に本発明によるコンパウンドを充填し、ボンド磁石を形成する。ロータ３００の複数の磁石スロット３０４に対向するようにＳ極とＮ極とが交互に配置されたステータ（不図示）と組み合わせることによってＩＰＭ型モータが得られる。
【０２４４】
ボンド磁石の成形は、種々の方法で実行することができる。例えば、熱硬化性樹脂を含むコンパウンドを用いる場合、スロット内圧縮成形法（例えば特開昭６３−９８１０８号公報参照）を採用することができる。また、熱可塑性樹脂を含むコンパウンドを用いる場合には、押出成形法や射出成形法を採用することができる。何れの成形方法を採用する場合においても、本発明によるコンパウンドは充填性に優れるので、スロット３０４ａおよび３０４ｂ内に確実に充填でき、且つ、機械特性や耐熱性が優れるとともに、従来と同等以上の磁気特性を有するボンド磁石を形成することができる。従って、従来よりも高性能で高信頼性の小型ＩＰＭ型モータを作製することが可能になる。
【０２４５】
本発明によるコンパウンドは、図２０（ａ）に示す角度センサ（ロータリーエンコーダ）４００が有するボンド磁石の形成に好適に用いられる。
【０２４６】
図２０（ａ）に示したロータリエンコーダ４１１は、回転軸４１３と、回転軸４１３に連結された回転ドラム４１６と、回転ドラム４１６の外周面に接合された複数のボンド磁石４１５とを有するロータ４１４と、ロータ４１４の外周面に離間して配置された検出器４１７とを備えている。検出器４１７は、ロータ４１４からの磁束の変化を検出できるものであれば、特に限定されず、例えば、ホール素子、磁気抵抗素子、磁気インピーダンス効果素子を用いることができる。また、回転軸４１３はモータ４１２に連結されている。検出器１７は、図示しない計測部に接続されている。
【０２４７】
本発明によるコンパウンドを用いて形成されたボンド磁石４１５は、例えば、図２０（ｂ）に示したような円柱状のものであり、回転ドラム４１６の外周面に沿ってＮ極とＳ極とが交互に配置されている。ボンド磁石４１５と回転ドラム４１６との接合は、例えば接着剤等によって行われている。回転ドラム４１６は、例えば、金属材料を用いて形成され、磁性材料でなくてもよい。
【０２４８】
このロータリエンコーダ４００は、以下のように動作する。モータ４１２の回転軸４１３が回転すると、その回転に応じてロータ４１４が回転する。このとき、ロータ４１４の外周面に配置されたボンド磁石４１５が検出器４１７に形成する磁束の向きが、ロータ４１４の回転に従って変化する。検出器４１７はこの磁束の向きの変化に相当する出力信号（電圧の変化量または電流の変化量等）を生成し、計測部（不図示）に出力する。このようにして、モータ４１２の回転量（角度）が計測される。
【０２４９】
本発明によるコンパウンドは充填性（成形性、流動性）に優れ、従来と同等以上の磁気特性を有し、且つ、従来よりも機械特性や耐熱性に優れるボンド磁石を形成することができるので、小型で高性能で信頼性の高い角度センサを作製することが可能になる。
【０２５０】
さらに、本発明によるコンパウンドは、図２１（ａ）および（ｂ）を参照しながら説明する磁気ロール用のボンド磁石の形成に好適に用いられる。
【０２５１】
図２１（ａ）は、電子写真用のプロセスカートリッジ５０１の構造を模式的に示す断面図である。カートリッジ５０１は、矢印Ａ方向に回転駆動される感光ドラム５１０と、感光ドラム５１０を帯電するための帯電ローラ５０２と、現像装置５１１と、クリーニング装置５１２とを一体に有している。
【０２５２】
現像装置５１１は、トナー５１３を収容する現像容器５０９を備え、現像容器５０９の開口部には感光ドラム５１０に対向するように現像スリーブ５０６が回転可能に配設されている。また、現像装置５１１は弾性ブレード５１４を備えており、弾性ブレード５１４は現像スリーブ５０６に当接し、現像スリーブ５０６により担持搬送されるトナー５１３の層厚を規制する。
【０２５３】
図２１（ｂ）は、プロセスカートリッジ５０１が有する現像装置５１１の構成を模式的に示す断面図である。
【０２５４】
現像スリーブ５０６は非磁性材料で形成されており、軸受を介して現像容器５０９に回転可能に支持されている。現像スリーブ（例えば直径１０ｍｍ）５０６内には磁気ロール（例えば直径８．５ｍｍ）５０７が配設されている。磁気ロール５０７の軸部５０７ａには切欠き５０７ａ−１が形成されており、切欠き５０７ａ−１が現像容器５０９に係合することによって磁気ロール５０７が固定されている。磁気ロール５０７は、感光ドラム５１０と対向する位置に現像極Ｓ１を有し、その他の位置にはＳ２極、Ｎ１極およびＮ２極を有している。
【０２５５】
磁石５０８は、現像スリーブ５０６を包囲するように配置されており、現像スリーブ５０６との間隙ｇに磁気カーテンを形成し、この磁気カーテンによって間隙内にトナーを保持することによって、トナー漏れが防止される。
【０２５６】
磁気ロール５０７は本発明によるコンパウンドを用いて形成されているので、従来の磁石と同等以上の磁気特性を有し、且つ、機械特性や耐熱性にも優れる。従って、磁気ロール５０７や現像スリーブ５０６を従来よりも更に小型化するこが可能であるとともに、性能を向上することができる。本発明によるコンパウンドを用いて形成された磁気ロールは、複写機やレーザビームプリンタ内の現像装置や現像カートリッジにも適用できる。
【０２５７】
【実施例】
（実施例１）
下記の表７に示す組成を有し、総量が６００グラム（ｇ）となるように純度９９．５％以上のＢ、Ｆｅ、Ｔｉ、Ｎｄ、およびＮｂを秤量し、アルミナ製坩堝に投入した。その後、これらの合金原料を高周波加熱によって圧力７０ｋＰａのアルゴン（Ａｒ）雰囲気中で溶解し、合金溶湯を作製した。溶湯温度が１５００℃に到達した後、水冷した銅製鋳型上に鋳込み、平板状の合金を作製した。
【０２５８】
得られた合金を粉砕した後、２５ミリグラム（ｍｇ）の粉砕片を溶解し、Ａｒ気流中で示差熱量計（ＤＴＡ）を用い、冷却速度２０℃／分で合金溶湯の凝固過程を解析した。測定結果を表７に示す。
【０２５９】
【表７】

【０２６０】
ここで、試料Ｎｏ．１〜５は、Ｔｉに加えてＮｂを添加した試料であり、試料Ｎｏ．６〜７はＮｂを添加しなかった試料である。
【０２６１】
表７の最右欄には、各試料Ｎｏ．１〜７の合金溶湯について、合金溶湯の凝固過程を特徴付ける温度を記載している。「１ｓｔ」と表記されている温度は、合金溶湯を冷却する過程で最初の凝固が生じた温度（「液相線温度」）を示している。「２ｎｄ」と表記されている温度は、合金溶湯を冷却する過程で液相線温度より低い温度で次の凝固が生じた温度（「凝固点」）を示している。これらの温度は、具体的には、示差熱熱量計（ＤＴＡ）によって発熱ピークが観測された温度である。
【０２６２】
図８は、試料Ｎｏ．２（Ｎｂ添加）および試料Ｎｏ．６（Ｎｂ無添加）のＤＴＡを示すグラフである。図８から明らかなように、試料Ｎｏ．２の場合、試料Ｎｏ．６に比較して、冷却過程で生じる最初の発熱ピークの温度、すなわち液相線温度（「１ｓｔ」）が６０℃以上も低下している。
【０２６３】
この最初の発熱ピークは、ＴｉＢ₂などのチタンとホウ素の化合物相が析出することに起因している可能性がある。本実施例では、Ｔｉおよびホウ素を従来よりも高濃度に添加しているため、チタンとホウ素の化合物（高融点）が形成されやすく、その析出温度は高いと推定される。Ｔｉを添加しない従来の組成系（Ｆｅ₃Ｂ／Ｎｄ₂Ｆｅ₁₄Ｂ系）では、合金溶湯の液相線温度は１２００℃程度以下であった。本発明では、ＴｉとともにＮｂが添加されることにより、このような化合物の析出温度が下がり、合金溶湯の液相線温度が低下したものと考えられる。
【０２６４】
試料Ｎｏ．６（比較例）の合金を用いる場合は、１３５０℃程度の高い出湯温度でストリップキャスティングを実行する必要があるが、試料Ｎｏ．２（実施例）の合金を用いる場合は、出湯温度を例えば約１２５０℃程度に設定することが可能である。このように出湯温度が低減されると、溶湯の冷却過程で早くに析出するＲ₂Ｆｅ₁₄Ｂ型化合物やＴｉＢ₂の粗大化が抑制され、磁石特性が向上する。
【０２６５】
（実施例２）
下記の表２に示す組成を有し、総量が６００ｇとなるように純度９９．５％以上のＢ、Ｆｅ、Ｔｉ、Ｎｄ、およびＣを秤量し、アルミナ製坩堝に投入した。その後、これらの合金原料を高周波加熱によって圧力７０ｋＰａのアルゴン（Ａｒ）雰囲気中で溶解し、合金溶湯を作製した。溶湯温度が１５００℃に到達した後、水冷した銅製鋳型上に鋳込み、平板状の合金を作製した。
【０２６６】
得られた合金を粉砕した後、２５ミリグラム（ｍｇ）の粉砕片を溶解し、Ａｒ気流中で示差熱熱量計（ＤＴＡ）を用い、冷却速度２０℃／分で合金溶湯の凝固過程を解析した。測定結果を表８に示す。
【０２６７】
【表８】

【０２６８】
ここで、試料Ｎｏ．８〜１３は、ＴｉとともにＣを添加した試料であり、試料Ｎｏ．１４〜１５はＣを添加しなかった試料である。
【０２６９】
表８の最右欄には、各試料Ｎｏ．８〜１５の合金溶湯について、合金溶湯の凝固過程を特徴付ける温度を記載している。「１ｓｔ」と表記されている温度は、合金溶湯を冷却する過程で最初の凝固が生じた温度（「液相線温度」）を示している。「２ｎｄ」と表記されている温度は、合金溶湯を冷却する過程で液相線温度より低い温度で次の凝固が生じた温度（「凝固点」）を示している。これらの温度は、具体的には、示差熱熱量計（ＤＴＡ）によって発熱ピークが観測された温度である。
【０２７０】
図９は、試料Ｎｏ．８（Ｃ添加）および試料Ｎｏ．１４（Ｃ無添加）のＤＴＡを示すグラフである。図９から明らかなように、試料Ｎｏ．８の場合、試料Ｎｏ．１４に比較して、冷却過程で生じる最初の発熱ピークの温度、すなわち液相線温度（「１ｓｔ」）が４０℃程度も低下している。
【０２７１】
この最初の発熱ピークは、ＴｉＢ₂などのチタンとホウ素の化合物相が析出することに起因している可能性がある。本実施例では、Ｔｉおよびホウ素を従来よりも高濃度に添加しているため、チタンとホウ素の化合物（高融点）が形成されやすく、その析出温度は高いと推定される。Ｔｉを添加しない従来の組成系（Ｆｅ₃Ｂ／Ｎｄ₂Ｆｅ₁₄Ｂ系）では、合金溶湯の液相線温度は１２００℃程度以下であった。本発明の実施例では、ＴｉとともにＣが添加されていたことにより、このような化合物の析出温度が下がり、合金溶湯の液相線温度が低下したものと考えられる。
【０２７２】
試料Ｎｏ．１４の合金を用いる場合は、１３５０℃程度の高い出湯温度でストリップキャスティングなどを実行する必要があるが、試料Ｎｏ．８（実施例）の合金を用いる場合は、出湯温度を例えば約１３００℃程度に設定することが可能である。このように出湯温度が低減されると、溶湯の冷却過程で早くに析出するＲ₂Ｆｅ₁₄Ｂ型化合物やＴｉＢ₂の粗大化が抑制され、磁石特性が向上する。
【０２７３】
次に、表８に示す組成を有し、総量が１５ｇとなるように純度９９．５％以上のＢ、Ｆｅ、Ｔｉ、Ｎｄ、およびＣを秤量し、底部に直径０．８ｍｍのオリフィスを有する石英坩堝に投入した。その後、これらの合金原料を高周波加熱によって圧力１．３３〜４７．９２ｋＰａのＡｒ雰囲気中で溶解し、合金溶湯を作製した。溶湯温度が１３５０℃に到達した後、湯面をＡｒガスで加圧し、オリフィスから溶湯を０．７ｍｍ下方に位置する冷却ロールの外周面へ滴下した。冷却ロールは純銅製であり、外周面速度が１５ｍ／秒となるように回転させていた。このような冷却ロールとの接触により、合金溶湯は急冷され、凝固した。こうして、幅２〜３ｍｍ、厚さ２０〜５０μmの連続した急冷凝固合金の薄帯が得られた。図１０は、試料Ｎｏ．８および試料Ｎｏ．１４のＸＲＤのパターンを示すグラフである。図１０から明らかなように、試料Ｎｏ．８の場合、非晶質が大部分を占めているのに対して、試料Ｎｏ．１４では、結晶組織の割合が多い。
【０２７４】
この急冷凝固合金薄帯をＡｒ雰囲気中において６００〜８００℃の熱処理温度範囲で６〜８分間保持し、その後、室温まで冷却した。この後、ＶＳＭを用いて急冷合金薄帯（長さ３〜５ｍｍ）の磁気特性を評価した。測定結果を表９に示す。
【０２７５】
【表９】

【０２７６】
次に、表８の試料Ｎｏ．１３と同一の組成を有する原料合金を用意し、図３に示すようなストリップキャスト装置を用いて、急冷合金を作製した。具体的には、総量が１０ｋｇとなるように純度９９．５％以上のＢ、Ｆｅ、Ｔｉ、Ｎｄ、およびＣを秤量し、溶解槽に投入した。その後、これらの合金原料を高周波加熱によって圧力３０ｋＰａのＡｒ雰囲気中で溶解し、合金溶湯を作製した。溶湯温度が１３５０℃に到達した後、溶湯をシュートに流した。溶湯はシュート上をスムーズに流れ、冷却ロールにより冷却された。冷却ロールの表面周速度は１２ｍ／秒とした。
【０２７７】
こうして得られた急冷合金（平均厚さ：８０μm程度）をＡｒ雰囲気中において７４０℃の熱処理温度範囲で６〜８分間保持し、その後、室温まで冷却した。その後、ＶＳＭを用いて急冷合金の磁気特性を評価した。
【０２７８】
測定結果は、残留磁束密度Ｂ_rが０．７９Ｔ、保磁力Ｈ_cJが１０９０ｋＡ／ｍ、最大磁気エネルギ積（ＢＨ）_maxが１０２ｋＪ／ｍ³であった。この磁気特性を表９に示す試料Ｎｏ．８の磁気特性と比較すると、ほぼ同等の特性が得られたことがわかる。
【０２７９】
次に、Ｃが（Ｂ＋Ｃ）の合計に占める割合（原子比率ｐ）が０．２５以下の試料と、ｐが０．２５を超える試料について、ＸＲＤ及び減磁曲線を測定した。
【０２８０】
図１１は、Ｎｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₄（実施例：ｐ＝０．１）及びＮｄ₉Ｆｅ₇₃Ｂ₇Ｃ₇Ｔｉ₄（比較例：ｐ＝０．５）の熱処理前におけるＸＲＤパターンを示している。これらの試料は、組成は異なるが、いずれも前述した実施例と同様にして作製された。なお、図１２は、Ｎｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₄（実施例）及びＮｄ₉Ｆｅ₇₃Ｂ₇Ｃ₇Ｔｉ₄（比較例）の減磁曲線を示している。
【０２８１】
Ｃの比率ｐが０．２５を超えて０．５に達する場合、図１１に示されるように、Ｔｉ−Ｃの回折ピークが顕著に観察される。このようにＣが多過ぎると、急冷合金中にＴｉ−Ｃ相が多く析出するため、熱処理後の構成相比率が所望範囲からずれ、図１２に示すように減磁曲線の角形性が悪くなる。Ｃが（Ｂ＋Ｃ）の合計に占める割合（原子比率ｐ）が０．２５以下であれば、このような問題は生じなかった。
【０２８２】
（実施例３）
本実施例では、図３に示すストリップキャスト装置を用いた。
【０２８３】
まず、原子比率でＮｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₃Ｎｂ₁の組成を有するように、純度９９.５%以上のＢ、Ｃ、Ｆｅ、Ｎｂ、Ｔｉ、およびＮｄの金属を用いて総量が５ｋｇとなるように秤量した。これらの金属をアルミナ製坩堝内に投入し、圧力３５ｋＰａのアルゴン雰囲気中で高周波加熱により溶解した。溶解温度は１３５０℃とした。
【０２８４】
溶解後、坩堝を傾転し、溶湯を多孔質セラミックス製のシュート上に供給し、冷却ロールの表面へ導いた。シュートの表面温度はヒータによって６００℃に保持した。また、シュート上において溶湯がロールへ向かってスムーズに流れるように、シュートを水平方向に対して２０°（＝角度α）だけ傾けた。また、溶湯は、ロールの直上部から坩堝の位置へ４０°（＝角度β）だけ傾斜した位置に注がれるようにシュートを配置した。なお、本実施例におけるシュートは、図４に示すように、坩堝から受けた溶湯の流れを２条に分けてロールへ供給するための溶湯ガイドを有している。
【０２８５】
冷却ロールは１４ｍ／秒の表面周速度で回転させた。坩堝の傾転角を調整することにより、シュート上を流れる溶湯の供給速度を、溶湯の１つの流れあたり１．５ｋｇ／分になるよう調整した。本実施例では、表面の中心線粗さＲａが５μｍの純銅製ロールを用いた。ロール温度の上昇はロール内部の水冷によって防止した。
【０２８６】
得られた急冷合金の組織をＣｕＫαの特性Ｘ線により調べたところ、Ｎｄ₂Ｆｅ₁₄Ｂの回折ピークとともに、Ｆｅ₂₃Ｂ₆およびα−Ｆｅが混在している急冷合金組織であることを確認した。
【０２８７】
図１３は、得られた急冷合金の粉末ＸＲＤを示し、図１４は振動型磁力計を用いて測定した急冷合金の減磁曲線を示す。図１３および図１４において、「ａｓ−ｃａｓｔ」と記載している曲線が急冷合金に関するものである。
【０２８８】
次に、急冷合金をパワーミルによって粉砕した。その後、アルゴンガスで流気し、炉内温度を７４０℃に保持したフープベルト式連続熱処理炉内に急冷合金粉末を供給して熱処理を施した。このとき給粉速度は３０ｇ／分に保持した。
【０２８９】
熱処理後における粉末ＸＲＤおよび減磁曲線も、それぞれ、図１３および図１４に示している。図１３および図１４において、熱処理後のデータは「ａｓ−ａｎｎｅａｌｅｄ」と記載された曲線で示されている。熱処理後の磁気特性を以下の表１０に示す。
【０２９０】
【表１０】

【０２９１】
図１４および表１０からわかるように、本実施例における鉄基永久磁石は良好な磁気特性を発揮した。
【０２９２】
次に、熱処理後の微細金属組織を透過型電子顕微鏡（ＴＥＭ）にて観測した。その結果、熱処理後の組織内には、平均粒径４０ｎｍ程度の結晶粒と、その粒界に１０ｎｍ程度の微細結晶粒とが存在していることがわかった。また、ＨＲＴＥＭ（高解像透過電子顕微鏡）による金属組織解析の結果、平均粒径４０ｎｍ程度の結晶粒はＮｄ₂Ｆｅ₁₄Ｂであり、その粒界にはＦｅ₂₃Ｂ₆またはＦｅ₃Ｂの鉄基硼化物が存在していることを確認した。
【０２９３】
（実施例４）
本実施例でも、図３に示すストリップキャスト装置を用いた。
【０２９４】
まず、原子比率でＮｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₃Ｎｂ₁の組成を有するように、純度９９.５%以上のＢ、Ｃ、Ｆｅ、Ｎｂ、Ｔｉ、およびＮｄの金属を用いて総量が５ｋｇとなるように秤量した。これらの金属をアルミナ製坩堝内に投入し、圧力３５ｋＰａのアルゴン雰囲気中で高周波加熱により溶解した。溶解温度は１３５０℃とした。
【０２９５】
溶解後、坩堝を傾転し、溶湯を多孔質セラミックス製のシュート上に供給し、冷却ロールの表面へ導いた。シュートの表面温度はヒータによって６００℃に保持した。また、シュート上において溶湯がロールへ向かってスムーズに流れるように、シュートを水平方向に対して２０°（＝角度α）だけ傾けた。また、溶湯は、ロールの直上部から坩堝の位置へ４０°（＝角度β）だけ傾斜した位置に注がれるようにシュートを配置した。本実施例でも、図４に示すシュートを用いた。
【０２９６】
本実施例では、表１１に示す表面周速度で冷却ロールを回転させた。また、坩堝の傾転角を調整することにより、シュート上を流れる溶湯の供給速度（１つの流れあたり）を表１１に示すように調整した。溶湯の１つの流れの幅は１０ｍｍとして、ロール周速度および溶湯供給速度が急冷に与える影響を調べた。
【０２９７】
なお、本実施例でも、実施例３と同様に、表面の中心線粗さＲａが５μｍの純銅製ロールを用いた。ロール温度の上昇はロール内部の水冷によって防止した。
【０２９８】
【表１１】

【０２９９】
表１１において、「○」は、安定して急冷合金を作製できた場合を示している。これに対し、「×」はスプラッシュが発生し、所望の組織を有する急冷合金を安定して得ることができなかった場合を示している。「△」は、安定した急冷合金の作製がしばしば観察されたものの、断続的にスプラッシュが発生した場合を示している。
【０３００】
表１１から、ロール表面周速度が１０ｍ／秒以上１８ｍ／秒以下の場合、１つの溶湯流れあたりの溶湯供給速度が１．０ｋｇ／分以上２．０ｋｇ／分以下で安定した急冷が実現していることがわかる。ロール表面周速度が速くなるほど、急冷合金薄帯は薄くなり、また、スプラッシュも発生しやすくなる。
【０３０１】
溶湯の１つの流れあたりの溶湯供給速度は、急冷合金薄帯の厚さにはさほど影響しないが、急冷合金薄帯の幅を変化させる。溶湯供給速度が大きいほど、急冷合金薄帯の幅が広くなる。
【０３０２】
急冷合金薄帯の厚さは、ロール表面周速度に依存して変化する。すなわち、ロール表面周速度が速いほど、急冷合金薄帯は薄くなる。例えば、ロール表面周速度が１０ｍ／秒のとき、急冷合金薄帯の平均厚さは１００μｍ程度であり、ロール表面周速度が２２ｍ／秒のとき、急冷合金薄帯の平均厚さは４５〜８０μｍ程度である。
【０３０３】
前述したように、急冷合金薄帯の厚さが厚いほど（例えば８０μｍを超える厚さを持つ場合）、急冷合金を粉砕することによって、等軸形状に近い形状の粉末粒子が得られやすい。アスペクト比が１に近い粒子が多く含まれる粉末を用いてボンド磁石を作製すれば、磁石特性に優れたボンド磁石を得ることができる。
【０３０４】
なお、ロールの表面周速度１４ｍ／秒、溶湯の１つの流れあたりの溶湯供給量１．３ｋｇ／分の条件で作製された急冷合金の組織をＣｕＫαの特性Ｘ線により調べた。その結果、Ｎｄ₂Ｆｅ₁₄Ｂの回折ピークとともに、Ｆｅ₂₃Ｂ₆およびα−Ｆｅが混在している急冷合金組織であることを確認した。
【０３０５】
【発明の効果】
本発明によれば、ストリップキャスト法を用い、Ｔｉを原料合金に添加した希土類合金溶湯の急冷を行なうことにより、磁石に必要な希土類元素の量を低減しながら保磁力および磁化が充分に高い優れた磁気特性を発揮する鉄基希土類磁石用原料合金を量産することができる。
【図面の簡単な説明】
【図１】本発明で好適に使用されるストリップキャスト装置の構成例を示す図である。
【図２】本発明により製造されるナノコンポジット磁石の組織を示す図である。
【図３】本発明で好適に使用されるストリップキャスト装置の他の構成例を示す図である。
【図４】ストリップキャスト装置で用いられる合金溶湯のシュート（案内手段）を示す斜視図である。
【図５】ストリップキャスト法に用いる冷却ロールの表面における中心線粗さＲａが合金溶湯の急冷に与える影響を示す図である。
【図６】メルトスピニング法に用いる冷却ロールの表面における中心線粗さＲａが合金溶湯の急冷に与える影響を示す図である。
【図７】ストリップキャスト法で形成された急冷合金の組織構造を示す断面図であり、（ａ）はＴｉを添加したＲ−Ｔ−Ｂ系合金の断面を示し、（ｂ）はＴｉを添加しない従来のＲ−Ｔ−Ｂ系合金の断面を示している。
【図８】試料Ｎｏ．２および試料Ｎｏ．６のＤＴＡを示すグラフである。
【図９】試料Ｎｏ．８および試料Ｎｏ．１４のＤＴＡを示すグラフである。
【図１０】結晶化熱処理前（ａｓ−ｃａｓｔ）の試料Ｎｏ．８および試料Ｎｏ．１４の粉末Ｘ線回折データを示すグラフである。
【図１１】Ｎｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₄（実施例：ｐ＝０．１）及びＮｄ₉Ｆｅ₇₃Ｂ₇Ｃ₇Ｔｉ₄（比較例：ｐ＝０．５）の熱処理前におけるＸＲＤパターンを示している。
【図１２】Ｎｄ₉Ｆｅ₇₃Ｂ_12.6Ｃ_1.4Ｔｉ₄（実施例）及びＮｄ₉Ｆｅ₇₃Ｂ₇Ｃ₇Ｔｉ₄（比較例）の減磁曲線を示している。
【図１３】本発明の実施例に関する粉末ＸＲＤのグラフである。「ａｓ−ｃａｓｔ」と記載している曲線が急冷合金に関するものであり、「ａｓ−ａｎｎｅａｌｅｄ」と記載している曲線が熱処理後における合金に関するものである。
【図１４】振動型磁力計を用いて測定した本発明の実施例に関する減磁曲線のグラフである。「ａｓ−ｃａｓｔ」と記載している曲線が急冷合金に関するものであり、「ａｓ−ａｎｎｅａｌｅｄ」と記載している曲線が熱処理後における合金に関するものである。
【図１５】本発明によるナノコンポジット磁粉および従来の急冷磁石粉末の加熱質量増加率を示すグラフである。
【図１６】粒度分布の異なるナノコンポジット磁粉を用いて形成されたボンド磁石成形体の密度を示すグラフである。
【図１７】本発明による実施形態の永久磁石回転子型を備えるステッピングモータ１００の構成を模式的に示す分解斜視図である。
【図１８】（ａ）〜（ｄ）は、本発明による実施形態のボンド磁石一体成形型のロータ２００およびその成形工程を示す図である。
【図１９】本発明による実施形態の磁石埋設型ロータ３００の構造を示す模式図である。
【図２０】（ａ）および（ｂ）は、本発明による実施形態のロータリーエンコーダ４１１の構造を模式的に示す図である。
【図２１】（ａ）および（ｂ）は、本発明による実施形態の磁気ロール５０７を備える電子写真用のプロセスカートリッジ５０１の構造を模式的に示す断面図である。
【符号の説明】
１ 溶解炉
２ 溶解炉の底部出口
３ 合金溶湯
４ 樋
５ シュート（溶湯の案内手段）
６ 合金溶湯のパドル
７ 冷却ロール
８ 急冷合金[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for manufacturing a permanent magnet suitably used for various motors and actuators, and particularly to an iron-based rare earth alloy magnet having a plurality of ferromagnetic phases and a method for manufacturing the same.
[0002]
[Prior art]
In recent years, further improvement in performance and reduction in size and weight have been required for home appliances, OA devices, electrical components, and the like. Therefore, for permanent magnets used in these devices, it is required to maximize the performance-to-weight ratio of the entire magnetic circuit. For example, residual magnetic flux density B _r Is required to use a permanent magnet of 0.5T (Tesla) or more. However, depending on the conventional relatively inexpensive hard ferrite magnet, the residual magnetic flux density B _r Cannot be 0.5T or more.
[0003]
Currently, high residual magnetic flux density B of 0.5T or more _r As permanent magnets having Sm-Co, Sm-Co magnets produced by powder metallurgy are known. Other than Sm—Co magnets, Nd—Fe—B sintered magnets produced by powder metallurgy and Nd—Fe—B quenching magnets produced by liquid quenching are high in residual magnetic flux density B. _r Can be demonstrated. The former Nd—Fe—B based sintered magnet is disclosed in, for example, Japanese Patent Laid-Open No. 59-46008, and the latter Nd—Fe—B based quenched magnet is disclosed in, for example, Japanese Patent Laid-Open No. 60-9852. Has been.
[0004]
[Problems to be solved by the invention]
However, Sm-Co magnets have the disadvantage of high magnet prices because both Sm and Co as raw materials are expensive.
[0005]
Nd-Fe-B magnets contain inexpensive Fe as a main component (about 60% to 70% by weight of the total), and are therefore less expensive than Sm-Co magnets. There is a problem that the cost required is high. One of the reasons for the high manufacturing process cost is that a large-scale facility and a large number of processes are required for the separation and purification of Nd, in which the content is about 10 atomic% to 15 atomic%, and the reduction reaction. In addition, in the case of the powder metallurgy method, the number of manufacturing steps is inevitably increased.
[0006]
On the other hand, an Nd—Fe—B type quenching magnet manufactured by a liquid quenching method can be obtained by a relatively simple process such as a melting process → a liquid cooling process → a heat treatment process. There is an advantage that the process cost is lower than that of the system magnet. However, in the case of the liquid quenching method, in order to obtain a bulk permanent magnet, it is necessary to mix a magnet powder produced from a quenching alloy with a resin to form a bonded magnet. The filling rate (volume ratio) is at most about 80%. Moreover, the quenched alloy produced by the liquid quenching method is magnetically isotropic.
[0007]
For the above reasons, the Nd—Fe—B type quenching magnet manufactured by using the liquid quenching method has a higher B than the anisotropic Nd—Fe—B type sintered magnet manufactured by the powder metallurgy method. _r Has the problem of low.
[0008]
The method for improving the characteristics of the Nd—Fe—B quenching magnet is selected from the group consisting of Zr, Nb, Mo, Hf, Ta, and W as described in Japanese Patent Laid-Open No. 1-7502. It is effective to add at least one element combined with at least one element selected from the group consisting of Ti, V, and Cr. By adding such elements, the coercive force H _cJ And corrosion resistance are improved, but residual magnetic flux density B _r There is no known effective method for improving the thickness other than increasing the density of the bonded magnet. In addition, when the Nd-Fe-B quenching magnet contains a rare earth element of 6 atomic% or more, according to many prior arts, in order to increase the quenching rate of the melt, the melt is injected to the cooling roll through the nozzle. A melt spinning method is used.
[0009]
In the case of an Nd-Fe-B quenching magnet, the composition of the rare earth element is relatively low, that is, Nd _3.8 Fe _77.2 B ₁₉ (Atom%) near composition, Fe _Three Magnet materials having a B-type compound as the main phase have been proposed (R. Coehoorn et al., J. de Phys, C8, 1998, pages 669-670). This permanent magnet material is obtained by applying a crystallization heat treatment to an amorphous alloy produced by a liquid quenching method, thereby producing Fe which is soft magnetic. _Three Nd which is B phase and hard magnetic ₂ Fe ₁₄ It has a metastable structure formed from a fine crystal aggregate in which the B phase is mixed, and is called a “nanocomposite magnet”. For such a nanocomposite magnet, a high residual magnetic flux density B of 1T or more _r Has been reported to have a coercive force H _cJ Is relatively low at 160 kA / m to 240 kA / m. Therefore, the use of this permanent magnet material is limited to applications where the operating point of the magnet is 1 or more.
[0010]
In addition, attempts have been made to improve various magnetic elements by adding various metal elements to the raw material alloy of the nanocomposite magnet (Japanese Patent Laid-Open No. 3-261104, Japanese Patent No. 2727505, Japanese Patent No. 2727506, International International Publication No. WO003 / 03403, WCChan, et.al. "THE EFFECTS OF REFRACTORY METALS ON THE MAGNETIC PROPERTIES OFα-Fe / R ₂ Fe ₁₄ B-TYPE NANOCOMPOSITES ", IEEE, Trans. Magn. No. 5, INTERMAG. 99, Kyongiu, Korea pp.3265-3267, 1999). This is because a coercive force of a size that can be practically used in a nanocomposite magnet has not been obtained, and sufficient magnetic properties cannot be expressed in actual use.
[0011]
The present invention has been made in view of the above circumstances, and its object is to provide a residual magnetic flux density B. _r High coercive force that can withstand practical use while maintaining ≧ 0.8T (for example, H _cJ An object of the present invention is to provide a method for producing a permanent magnet, which can produce an iron-based alloy magnet having excellent magnetic properties satisfying ≧ 600 kA / m) at low cost.
[0012]
[Means for Solving the Problems]
The method for producing an iron-based rare earth magnet raw material alloy according to the present invention has a composition formula of (Fe _1-m T _m ) _100-xyzn (B _1-p C _p ) _x R _y Ti _z M _n (T is one or more elements selected from the group consisting of Co and Ni, R is one or more elements selected from the group consisting of Y (yttrium) and rare earth metals, M is Al, Si, V, One or more elements selected from the group consisting of Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb), and a composition ratio ( Atomic ratio) x, y, z, m, n, and p are 10 <x ≦ 25 atomic%, 7 ≦ y <10 atomic%, 0.5 ≦ z ≦ 12 atomic%, 0 ≦ m ≦ 0, respectively. .5, a step of preparing a molten iron-based rare earth alloy satisfying 0 ≦ n ≦ 10 atomic% and 0 ≦ p ≦ 0.25, and a molten metal of the alloy, the guide surface is 1 in the horizontal direction. Supplied on guide means forming an angle of ˜80 °, the alloy in the contact area with the cooling roll And moving the hot water, the molten alloy was quenched by the chill roller, R ₂ Fe ₁₄ And a cooling step for producing a quenched alloy containing a B-type compound phase.
[0013]
In a preferred embodiment, the cooling step includes a step of adjusting the flow width of the molten alloy to a predetermined size along the axial direction of the cooling roll by the guide means.
[0014]
In a preferred embodiment, the quenched alloy is produced in a reduced-pressure atmosphere gas.
[0015]
In a preferred embodiment, the pressure of the atmospheric gas is adjusted to a pressure of 0.13 kPa to 100 kPa.
[0016]
In a preferred embodiment, in the cooling step, the R ₂ Fe ₁₄ The ratio of the B-type compound phase is 60% by volume or more of the quenched alloy.
[0017]
In a preferred embodiment, in the cooling step, a rotational peripheral speed of the surface of the cooling roll is adjusted to a range of 5 m / second to 26 m / second, and a supply speed per unit width of the molten alloy is 3 kg / minute / cm or less. And
[0018]
In a preferred embodiment, at least R ₂ Fe ₁₄ Forming a structure containing three or more crystal phases including a B-type compound phase, an α-Fe phase, and a ferromagnetic iron-based boride phase; ₂ Fe ₁₄ Including a step of setting an average crystal grain size of the B-type compound phase to 20 nm or more and 200 nm or less, and an average crystal grain size of the α-Fe phase and the boride phase of 1 nm to 50 nm.
[0019]
In a preferred embodiment, the ferromagnetic iron-based boride phase is R ₂ Fe ₁₄ It exists at the grain boundary or subgrain boundary of the B-type compound phase.
[0020]
In a preferred embodiment, the microstructure is formed by performing a crystallization heat treatment on the quenched alloy.
[0021]
In a preferred embodiment, the crystallization heat treatment includes holding the quenched alloy at a temperature of 550 ° C. or higher and 850 ° C. or lower for 30 seconds or longer.
[0022]
In a preferred embodiment, the method includes a step of pulverizing the quenched alloy before the crystallization heat treatment.
[0023]
In a preferred embodiment, the iron-based boride is Fe. _Three B and / or Fe _{twenty three} B ₆ Is included.
[0024]
In a preferred embodiment, the element M necessarily contains Nb.
[0025]
In a preferred embodiment, the liquidus temperature of the molten metal is 10 ° C. or more lower than that of an iron-based rare earth magnet raw material alloy having substantially the same composition except that Nb is not substantially contained.
[0026]
In a preferred embodiment, the Nb content is 0.1% or more and 3% or less of the whole in terms of atomic ratio.
[0027]
In a preferred embodiment, the composition ratio p of C in the composition formula satisfies a relationship of 0.01 ≦ p ≦ 0.25.
[0028]
In a preferred embodiment, the kinematic viscosity of the molten alloy before the supply to the guide means is 5 × 10. ^-6 (M ² / Sec) or less.
[0029]
In a preferred embodiment, the solidification temperature of the compound phase first precipitated in the solidification process of the molten alloy is reduced by 5 ° C. or more compared to the case where the composition ratio p is zero.
[0030]
In a preferred embodiment, in the cooling step, the compound phase first precipitated in the solidification process of the molten alloy is a titanium boride compound.
[0031]
In a preferred embodiment, the cooling step is performed by rotating a cooling roll having a surface centerline roughness Ra of 20 μm or less at a surface peripheral speed of 10 m / sec or more.
[0032]
In a preferred embodiment, in the cooling step, a molten metal quenching treatment rate per one flow of the molten alloy that is rapidly cooled by the cooling roll is adjusted within a range of 0.7 kg / min or more and less than 4 kg / min.
[0033]
In a preferred embodiment, in the cooling step, the width of one flow of the molten alloy is adjusted to 5 mm or more and less than 20 mm by the guide means.
[0034]
In a preferred embodiment, the kinematic viscosity of the molten alloy is 5 × 10. ^-6 m ² Adjust to less than / sec.
[0035]
In a preferred embodiment, the molten alloy has a kinematic viscosity of 5 × 10 5. ^-6 m ² The surface temperature of the guiding means is kept at 300 ° C. or higher so as not to exceed / sec.
[0036]
In a preferred embodiment, the thickness of the quenched alloy is 50 μm or more and 200 μm or less.
[0037]
In a preferred embodiment, the guiding means is Al. ₂ O _Three Is made of a material containing 80% by volume or more.
[0038]
In a preferred embodiment, the cooling roll uses a base material formed from a material having a thermal conductivity of 50 W / m / K or more.
[0039]
In a preferred embodiment, the cooling roll has a substrate formed of carbon steel, tungsten, iron, copper, molybdenum, beryllium, or a copper-based alloy.
[0040]
In a preferred embodiment, the surface of the base material of the cooling roll is plated with chromium, nickel, or a combination thereof.
[0041]
The method for producing an iron-based permanent magnet according to the present invention includes a step of preparing a quenched alloy produced by the above-described production method, and a step of performing a heat treatment on the quenched alloy.
[0042]
The manufacturing method of the bonded magnet by this invention includes the process of preparing the powder of the alloy powder produced by one of the said manufacturing methods, and the process of producing a bonded magnet using the said powder.
[0043]
The quenched alloy according to the present invention has a composition formula of (Fe _1-m T _m ) _100-xyzn Q _x R _y Ti _z M _n (T is one or more elements selected from the group consisting of Co and Ni, Q is one or more elements selected from the group consisting of B and C, R is a rare earth metal element, M is Al, Si, V, Cr, Mn, Cu, Ga, Zr, Nb, Mo, Hf, Ta, W, Pt, Pb, Au, and one or more elements selected from the group consisting of Ag and Ag), and a composition ratio (atomic ratio) ) X, y, z, m, and n are 10 <x ≦ 20 atomic%, 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and A quenched alloy satisfying 0 ≦ n ≦ 5 atomic%, having a thickness in the range of 50 μm to 200 μm, and having a crystal structure formed on two end faces perpendicular to the thickness direction To do.
[0044]
In a preferred embodiment, the crystal structure includes a ferromagnetic boride phase having an average particle diameter of 1 nm or more and 50 nm or less, and an R having an average particle diameter of 20 nm or more and 200 nm or less. ₂ Fe ₁₄ B-type compound phase.
[0045]
In a preferred embodiment, an amorphous portion exists in a region sandwiched between the crystal structures on both end faces.
[0046]
In a preferred embodiment, the thickness is 80 μm or more.
[0047]
The quenched alloy according to the present invention has a composition formula of (Fe _1-m T _m ) _100-xyzn Q _x R _y Ti _z M _n (T is one or more elements selected from the group consisting of Co and Ni, Q is one or more elements selected from the group consisting of B and C, R is a rare earth metal element, M is Al, Si, One or more elements selected from the group consisting of V, Cr, Mn, Cu, Ga, Zr, Nb, Mo, Hf, Ta, W, Pt, Pb, Au, and Ag), and a composition ratio x, y , Z, m, and n are 10 <x ≦ 20 atomic%, 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦, respectively. It is a quenched alloy that satisfies 5 atomic%, has a thickness in the range of 60 μm to 150 μm, and a recoil permeability of 1.1 to 2 inclusive.
[0048]
The magnetic powder according to the present invention has the composition formula (Fe _1-m T _m ) _100-xyzn Q _x R _y Ti _z M _n (T is one or more elements selected from the group consisting of Co and Ni, Q is one or more elements selected from the group consisting of B and C, R is a rare earth metal element, M is Al, Si, One or more elements selected from the group consisting of V, Cr, Mn, Cu, Ga, Zr, Nb, Mo, Hf, Ta, W, Pt, Pb, Au, and Ag), and a composition ratio x, y , Z, m, and n are 10 <x ≦ 20 atomic%, 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦, respectively. Magnetic powder satisfying 5 atomic%, having an average particle diameter of 60 μm or more and 110 μm or less, a ratio of a major axis size to a minor axis size of 0.3 or more and 1 or less, _cJ Is 600 kA / m or more.
[0049]
DETAILED DESCRIPTION OF THE INVENTION
The method for producing a permanent magnet according to the present invention is a method of cooling an iron-based alloy melt containing Fe, B, R (one or more rare earth metal elements including Y) and Ti in a reduced-pressure atmosphere by a strip casting method. As a result, fine R ₂ Fe ₁₄ A quenched alloy containing a B-type compound phase is prepared. Then, if necessary, a heat treatment is performed on the quenched alloy to crystallize the amorphous material remaining in the quenched alloy.
[0050]
The strip casting method is a method for producing a ribbon of a quenched alloy by bringing a molten alloy into contact with the surface of a cooling roll and cooling the molten alloy. In the present invention, the molten alloy is rapidly cooled and solidified by a cooling roll rotating at a higher speed than the conventional strip casting method. Compared with the melt spinning method, which uses a nozzle orifice to inject molten alloy onto the surface of the cooling roll, the strip casting method has a lower cooling rate but can produce a wide and relatively thick quenched alloy ribbon. Excellent in properties.
[0051]
According to the present invention, almost no soft magnetic α-Fe is precipitated in the quenched alloy, and fine R ₂ Fe ₁₄ Crystal structure with B-type compound phase or fine R ₂ Fe ₁₄ A structure in which a structure having a B-type compound phase and an amorphous phase are mixed is produced. As a result, R ₂ Fe ₁₄ High-performance composite type that suppresses the coarsening of the B-type compound phase, has an average particle size of 20 nm to 150 nm even after heat treatment, and has a soft magnetic phase such as an α-Fe phase finely dispersed A permanent magnet can be obtained. The fine soft magnetic phase is R ₂ Fe ₁₄ It exists in the grain boundary or subgrain boundary of the B-type compound phase, and exchange interaction is strengthened between constituent phases.
[0052]
Conventionally, a molten alloy having a composition similar to the composition targeted by the present invention (that is, a composition obtained by removing Ti from the composition of the present invention) is cooled and R ₂ Fe ₁₄ When an attempt is made to produce a quenched alloy containing a large amount of B-type compound phase, an alloy structure in which a large amount of α-Fe is precipitated is obtained. For this reason, there existed a problem that (alpha) -Fe coarsened by subsequent crystallization heat processing. When the soft magnetic phase such as α-Fe is coarsened, the magnet characteristics are greatly deteriorated, and a permanent magnet that can withstand practical use cannot be obtained.
[0053]
In particular, when the boron content is relatively high and the rare earth element R is relatively low (10 atomic% or less) as in the raw material alloy composition used in the present invention, the cooling rate of the molten alloy is sufficiently reduced according to the prior art. Let R ₂ Fe ₁₄ If an attempt is made to produce a rapidly solidified alloy in which the volume ratio of the B-type compound phase exceeds 60%, R ₂ Fe ₁₄ In addition to the B-type compound phase, a large amount of α-Fe or its precursor was precipitated, and the subsequent crystallization heat treatment caused the α-Fe phase to become coarse, resulting in a significant deterioration of the magnet characteristics.
[0054]
From the above, in order to increase the coercivity of nanocomposite magnets, conventionally, the melt spinning method was used to increase the cooling rate of the molten alloy so that most of the rapidly solidified alloy was occupied by the amorphous phase. Later, there was a common sense that it is preferable to form a uniformly refined structure from the amorphous phase by crystallization heat treatment. This is because, in order to obtain a nanocomposite magnet having an alloy structure in which a fine crystal phase is dispersed, it was thought that crystallization should be performed from an amorphous phase in a heat treatment process that is easy to control.
[0055]
For this reason, La having excellent amorphous forming ability is added to a raw material alloy, and a rapidly solidified alloy having an amorphous phase as a main phase is prepared by quenching a molten metal of the raw material alloy, and then Nd is subjected to crystallization heat treatment. ₂ Fe ₁₄ A technique has been reported in which both the B phase and the α-Fe phase are precipitated and grown, and each phase is as fine as several tens of nanometers (WCChan, et.al. "THE EFFECTS OF REFRACTORY METALS ON THE MAGNETIC PROPERTIES OF α-Fe / R ₂ Fe ₁₄ B-TYPE NANOCOMPOSITES ", IEEE, Trans. Magn. No. 5, INTERMAG. 99, Kyongiu, Korea pp.3265-3267, 1999). This paper describes the addition of trace amounts of refractory metal elements such as Ti (2 at. %) Improves the magnetic characteristics and increases the composition ratio of Nd, which is a rare earth element, from 19.5 at% to 11.0 at%. ₂ Fe ₁₄ It teaches that it is preferable to refine both the B phase and the α-Fe phase. The addition of the above refractory metal is a boride (R ₂ Fe _{twenty three} B _Three Or Fe _Three B) is suppressed, Nd ₂ Fe ₁₄ This is performed to produce a magnet composed of only two phases of a B phase and an α-Fe phase.
[0056]
The quenched alloy for the above-mentioned nanocomposite magnet is produced by a melt spinning method in which a molten alloy is injected onto the surface of a cooling roll that rotates at high speed using a nozzle. In the case of the melt spinning method, an extremely fast cooling rate can be obtained, which is suitable for producing an amorphous quenched alloy.
[0057]
In contrast, in the present invention, the molten alloy is cooled at a rate slower than the cooling rate in the conventional melt spinning method using the strip cast method, but γ-Fe ( (Precipitation of an α-Fe phase later) is suppressed, and further, coarsening of a soft magnetic phase such as an α-Fe phase in a crystallization heat treatment step is suppressed. As a result, fine R ₂ Fe ₁₄ A quenched alloy in which the B-type compound phase is uniformly dispersed can be produced.
[0058]
According to the present invention, mass production of permanent magnets having high magnetization (residual magnetic flux density) and coercive force and excellent demagnetization curve squareness while using a raw material alloy with a relatively small amount of rare earth elements (less than 10 at%). Can be manufactured at level.
[0059]
The increase in coercivity according to the invention is Nd ₂ Fe ₁₄ B phase is preferentially precipitated and grown in the cooling process, thereby Nd ₂ Fe ₁₄ This is achieved by increasing the volume ratio of the B phase but suppressing the coarsening of the soft magnetic phase. The increase in magnetization is due to the action of Ti to generate a boride phase such as a ferromagnetic iron-based boride from the boron-rich nonmagnetic amorphous phase present in the rapidly solidified alloy, and the nonmagnetic amorphous phase remaining after the crystallization heat treatment This is considered to be obtained by reducing the volume ratio.
[0060]
Hereinafter, the iron-based rare earth alloy magnet of the present invention will be described in more detail.
[0061]
First, the composition formula is (Fe _1-m T _m ) _100-xyzn (B _1-p C _p ) _x R _y Ti _z M _n Prepare a molten iron-based rare earth alloy expressed by Here, T is one or more elements selected from the group consisting of Co and Ni, R is one or more elements selected from the group consisting of Y (yttrium) and a rare earth metal, M is Al, Si, One or more elements selected from the group consisting of V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb. The composition ratios (atomic ratios) x, y, z, m, n, and p each satisfy the following relational expression.
[0062]
10 <x ≦ 25 atomic%
7 ≦ y <10 atomic%
0.5 ≦ z ≦ 12 atomic%
0 ≦ m ≦ 0.5
0 ≦ n ≦ 10 atomic% and
0 ≦ p ≦ 0.25
[0063]
Next, the above molten alloy is rapidly cooled in a reduced-pressure atmosphere gas by a strip casting method, and fine (for example, the average particle diameter is 150 nm or less) R ₂ Fe ₁₄ A cooling step for producing a quenched alloy containing a B-type compound phase in a volume ratio of 60% or more is performed.
[0064]
Then, if necessary, a crystallization heat treatment is performed on the quenched alloy, and R ₂ Fe ₁₄ A nanocomposite structure including a B-type compound phase and a ferromagnetic iron-based boride phase is formed. As the soft magnetic phase, in addition to the iron-based boride, a fine α-Fe phase may be included. In such an organization, R ₂ Fe ₁₄ The cooling conditions and crystallization heat treatment conditions of the molten alloy are adjusted so that the average crystal grain size of the B-type compound phase is 20 nm to 200 nm and the average crystal grain size of the boride phase and α-Fe phase is 1 nm to 50 nm. The
[0065]
According to the present invention, due to the action of the added Ti, in the cooling process of the molten alloy, R ₂ Fe ₁₄ Many B-type compound phases can be produced preferentially.
[0066]
R in the final magnet ₂ Fe ₁₄ The average crystal grain size of the B-type compound phase is larger than the average crystal grain size of the iron-based boride phase or the α-Fe phase. R, which is a hard magnetic phase ₂ Fe ₁₄ When the average size of the B-type compound phase is relatively large and the average size of the soft magnetic phase such as the α-Fe phase is sufficiently small, the constituent phases are effectively combined by exchange interaction, and the magnetization direction of the soft magnetic phase Is constrained by the hard magnetic phase, so that the entire alloy can exhibit excellent demagnetization curve squareness.
[0067]
In the present invention, by adjusting manufacturing conditions such as alloy composition, alloy cooling rate, and heat treatment temperature, R ₂ Fe ₁₄ It becomes possible to produce iron-based borides and α-Fe having a saturation magnetization equivalent to or higher than the saturation magnetization of the B-type compound phase. The iron-based boride produced is, for example, Fe _Three B (saturation magnetization 1.5T) or Fe _{twenty three} B ₆ (Saturation magnetization 1.6 T). Where R ₂ Fe ₁₄ The saturation magnetization of B is about 1.6T when R is Nd, and the saturation magnetization of α-Fe is 2.1T.
[0068]
In the production method of the present invention, the reason why the ferromagnetic iron-based boride as described above is likely to be generated is that R ₂ Fe ₁₄ When a solidified alloy in which the B-type compound phase occupies the majority, the amorphous phase present in the quenched alloy inevitably contains an excessive amount of boron. Therefore, this excess boron is combined with other elements in the crystallization heat treatment. This is thought to be due to easy precipitation and growth. However, when boron and other elements contained in the amorphous phase before heat treatment are combined to produce a compound having low magnetization, the magnetization of the entire magnet is lowered. The “amorphous phase” in this specification is not only a phase constituted only by a part in which the atomic arrangement is completely disordered, but also a crystallization precursor, a microcrystal (size: several nm or less), or an atom It shall also include phases that partially contain clusters. Specifically, a phase in which the crystal structure cannot be clearly identified by X-ray diffraction or transmission electron microscope observation will be widely referred to as an “amorphous phase”. A structure that can clearly identify the crystal structure by X-ray diffraction or transmission electron microscope observation is referred to as a “crystal phase”.
[0069]
According to the inventor's experiment, only when Ti is added, unlike the case of adding other types of metals such as V, Cr, Mn, Nb, and Mo, the magnetization does not decrease, but rather the magnetization is improved. I found out that Further, when Ti was added, the squareness of the demagnetization curve was particularly good as compared with the other additive elements described above. From these facts, it is considered that Ti plays an especially important role in suppressing the formation of a boride having a low magnetization. In particular, when boron and Ti are relatively small in the composition range of the raw material alloy used in the present invention, an iron-based boride phase having ferromagnetism is likely to precipitate by heat treatment. In this case, boron contained in the non-magnetic amorphous phase is incorporated into the iron-based boride. As a result, the volume ratio of the non-magnetic amorphous phase remaining after the crystallization heat treatment is decreased, and the ferromagnetic crystalline phase is increased. , Residual magnetic flux density B _r Will improve.
[0070]
Moreover, when Ti is added, the grain growth of α-Fe is suppressed and excellent hard magnetic properties are exhibited. And R ₂ Fe ₁₄ Ferromagnetic phases other than the B phase and the α-Fe phase are generated, thereby making it possible to form a structure including three or more types of ferromagnetic phases in the alloy. When a metal element such as Nb, V, or Cr is added instead of Ti, the α-Fe phase grows significantly in a relatively high temperature region where the α-Fe phase precipitates, and α-Fe As a result that the magnetization direction of the phase is not effectively constrained by exchange coupling with the hard magnetic phase, the squareness of the demagnetization curve is greatly reduced.
[0071]
When Nb, Mo, and W are added instead of Ti, if the heat treatment is performed in a relatively low temperature range where α-Fe does not precipitate, good hard magnetic properties with excellent demagnetization curve squareness can be obtained. It is possible. However, in alloys that have been heat treated at such temperatures, R ₂ Fe ₁₄ It is presumed that the B-type fine crystal phase is dispersed in the non-magnetic amorphous phase, and the composition of the nanocomposite magnet is not formed. Further, when heat treatment is performed at a higher temperature, an α-Fe phase is precipitated from the amorphous phase. Unlike the case where Ti is added, the α-Fe phase grows rapidly and becomes coarse after precipitation. For this reason, the magnetization direction of the α-Fe phase is not effectively constrained by exchange coupling with the hard magnetic phase, and the squareness of the demagnetization curve is greatly deteriorated.
[0072]
On the other hand, when V or Cr is added instead of Ti, these added metals are dissolved in Fe and are antiferromagnetically coupled, so that the magnetization is greatly reduced.
[0073]
On the other hand, when Ti is added, the kinetics of the precipitation / growth of the α-Fe phase is slow, and it takes time for the precipitation / growth. Therefore, Nd before the precipitation / growth of the α-Fe phase is completed. ₂ Fe ₁₄ It is considered that the precipitation and growth of the B phase starts. For this reason, before the α-Fe phase becomes coarse, Nd ₂ Fe ₁₄ The B phase grows greatly in a uniformly dispersed state.
[0074]
Only when Ti is added in this manner, the coarsening of the α-Fe phase can be appropriately suppressed, and a ferromagnetic iron-based boride can be formed. Furthermore, Ti plays an important role together with boron and carbon as an element that delays the crystallization of the Fe primary crystal (γ-Fe, which is subsequently transformed into α-Fe) during liquid quenching and facilitates the formation of a supercooled liquid. The cooling rate when quenching the molten alloy is 10 ² ℃ / sec ~ 10 ^Four Even if a relatively low value of about ° C./second is used, R does not precipitate coarse α-Fe. ₂ Fe ₁₄ Quenched alloy (R) containing 60% by volume or more of B-type crystal phase ₂ Fe ₁₄ It may be possible to produce an iron-based boride in addition to the B-type crystal phase.
[0075]
In the present invention, a strip casting method is used in which the molten metal is poured directly from the chute (guide means) onto the cooling roll without controlling the flow rate of the molten metal by the nozzle orifice. For this reason, productivity is high and manufacturing cost is low compared with the case of the melt spinning method using a nozzle orifice. Thus, in order to make an R—Fe—B rare earth alloy melt amorphous in a cooling rate range that can be achieved even by a strip casting method, it is usually necessary to add 10 atomic% or more of B (boron). When a large amount of B is added in this way, a non-magnetic amorphous phase with a high B concentration remains in the metal structure even after the crystallization heat treatment of the quenched alloy, and a homogeneous fine crystal structure is obtained. Absent. As a result, the volume ratio of the ferromagnetic phase is reduced, leading to a decrease in magnetization. However, when Ti is added as in the present invention, the above-described phenomenon is observed, so that an iron-based boride having a high magnetization is generated, and the magnetization is unexpectedly improved.
[0076]
[Reason for limiting composition]
When the total composition ratio x of B and C is 10 atomic% or less, the cooling rate during rapid cooling is 10 ² ℃ / sec ~ 10 ^Five If it is relatively slow at around ° C / second, R ₂ Fe ₁₄ It becomes difficult to produce a quenched alloy in which a B-type crystal phase and an amorphous phase are mixed, and a high coercive force cannot be obtained even if heat treatment is performed thereafter. Further, when the composition ratio x is 10 atomic% or less, iron-based borides exhibiting high magnetization are not generated. Since boron in the iron-based boride combines with Ti to form a stable compound, the more iron-based boride, the better the weather resistance. For this reason, x needs to exceed 10 atomic%. On the other hand, when the composition ratio x exceeds 25 atomic%, the volume ratio of the amorphous phase remaining after the crystallization heat treatment increases, and at the same time, the abundance ratio of α-Fe having the highest saturation magnetization in the constituent phases decreases. , Residual magnetic flux density B _r Will fall. From the above, it is preferable to set the composition ratio x so as to exceed 10 atomic% and not more than 25 atomic%. A more preferable range of the composition ratio x is more than 10 atomic% and 17 atomic% or less.
[0077]
The ratio p of C to the total of B and C is preferably in the range of 0 to 0.25 in terms of atomic ratio. In order to obtain the effect of C addition, the C ratio p is preferably 0.01 or more. If p is less than 0.01, the effect of C addition is hardly obtained. On the other hand, when p is too larger than 0.25, the amount of α-Fe phase generated increases, resulting in a problem that the magnetic properties deteriorate. The lower limit of the ratio p is preferably 0.02, and the upper limit of p is preferably 0.20 or less. The ratio p is more preferably 0.08 or more and 0.15 or less.
[0078]
R is one or more elements selected from the group of rare earth elements (including Y). When La or Ce is present, R ₂ Fe ₁₄ Since R (typically Nd) of the B phase is substituted with La or Ce, and the coercive force and the squareness deteriorate, it is preferable that La and Ce are not substantially contained. However, when a very small amount of La or Ce (0.5 atomic% or less) exists as an unavoidable impurity, there is no problem in terms of magnetic characteristics. Therefore, it can be said that La and Ce are not substantially contained when 0.5 or less atomic percent La and Ce are contained.
[0079]
More specifically, R preferably contains Pr or Nd as an essential element, and part of the essential element may be substituted with Dy and / or Tb. When the composition ratio y of R becomes less than 6 atomic% of the total, R required for the expression of coercive force ₂ Fe ₁₄ A compound phase having a B-type crystal structure does not sufficiently precipitate and has a high coercive force H _cJ You will not be able to get. Further, when the composition ratio y of R is 10 atomic% or more, the abundance of iron-based boride having ferromagnetism decreases, and the abundance of the B-rich nonmagnetic layer increases instead. Not formed and magnetization decreases. Therefore, the composition ratio y of the rare earth element R is preferably adjusted to a range of 6 atomic% to less than 10 atomic%, for example, 6 atomic% to 9.5 atomic%. A more preferable range of R is 7 atom% or more and 9.3 atom% or less, and a most preferable range of R is 8.3 atom% or more and 9.0 atom% or less.
[0080]
The addition of Ti exhibits the effect of precipitating and growing the hard magnetic phase earlier than the soft magnetic phase during the rapid cooling of the molten alloy, and also has a coercive force H _cJ And residual magnetic flux density B _r Contributes to the improvement of the squareness of the demagnetization curve and the maximum energy product (BH) _max To improve.
[0081]
When the Ti composition ratio z is less than 0.5 atomic% of the whole, the effect of Ti addition is not sufficiently exhibited. On the other hand, if the composition ratio z of Ti exceeds 12 atomic%, the volume ratio of the amorphous phase remaining after the crystallization heat treatment increases, so the residual magnetic flux density B _r It is easy to invite a decline. From the above, the composition ratio z of Ti is preferably in the range of 0.5 atomic% to 12 atomic%. The lower limit of the more preferable z range is 1.0 atomic%, and the upper limit of the more preferable z range is 6 atomic%. Furthermore, the upper limit of the preferable range of z is 5 atomic%.
[0082]
Further, the higher the composition ratio x of Q composed of C and / or B, the easier it is to form an amorphous phase containing excessive Q (for example, boron). Therefore, it is preferable to increase the Ti composition ratio z. Ti has a strong affinity for B and is concentrated at the grain boundaries of the hard magnetic phase. If the ratio of Ti to B is too high, Ti will not be at grain boundaries, but R ₂ Fe ₁₄ There is a possibility of entering the B compound and lowering the magnetization. On the other hand, if the ratio of Ti to B is too low, many magnetic B-rich amorphous phases will be generated. According to experiments, it is preferable to adjust the composition ratio so as to satisfy 0.05 ≦ z / x ≦ 0.4, and it is more preferable to satisfy 0.1 ≦ z / x ≦ 0.35. More preferably, 0.13 ≦ z / x ≦ 0.3.
[0083]
In order to obtain various effects, the metal element M may be added. M is one or more elements selected from the group consisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Hf, Ta, W, Pt, Pb, Au, and Ag It is.
[0084]
Fe occupies the remainder of the above-mentioned elements, but desired hard magnetic properties can be obtained even if a part of Fe is replaced with one or two transition metal elements (T) of Co and Ni. When the substitution amount of T for Fe exceeds 50%, a high residual magnetic flux density B of 0.7 T or more _r Cannot be obtained. For this reason, the substitution amount is preferably limited to a range of 0% to 50%. By replacing part of Fe with Co, the squareness of the demagnetization curve is improved and R ₂ Fe ₁₄ Since the Curie temperature of the B phase is increased, the heat resistance is improved. A preferable range of the amount of Fe substitution by Co is 0.5% or more and 40% or less.
[0085]
Next, preferred embodiments of the present invention will be described in detail with reference to the drawings.
[0086]
(Embodiment 1)
First, a first embodiment of the present invention will be described.
[0087]
In this embodiment, a rapidly solidified alloy is produced using the strip casting apparatus shown in FIG. In order to prevent the oxidation of the raw material alloy containing rare earth elements R and Fe which are easily oxidized, the quenching alloy is manufactured in an inert gas atmosphere. As the inert gas, a rare gas such as helium or argon or nitrogen can be used. Note that since nitrogen is relatively easy to react with the rare earth element R, it is preferable to use a rare gas such as helium or argon.
[0088]
The strip casting apparatus of FIG. 1 is disposed in a chamber (not shown) that can be reduced in pressure in an inert gas atmosphere. The strip casting apparatus includes a melting furnace 1 for melting the alloy raw material, a cooling roll 7 for rapidly cooling and solidifying the molten alloy 3 supplied from the bottom outlet 2 of the melting furnace 1, and a cooling roll from the melting furnace 1. 7 includes a gutter 4 and a chute (guide means) 5 for guiding the molten metal 3 to a melter 7, and a scraper gas ejector 9 that solidifies and easily peels the ribbon-shaped alloy 8 from the cooling roll 7.
[0089]
The melting furnace 1 can supply the molten metal 3 produced by melting the alloy raw material to the chute 5 at a substantially constant supply amount. This supply amount can be arbitrarily adjusted by controlling the operation of tilting the melting furnace 1. Note that the rod 4 is not essential, and the molten alloy 3 discharged from the melting furnace 1 may be directly supplied to the chute 5.
[0090]
The outer peripheral surface of the cooling roll 7 is formed from a material having good thermal conductivity such as copper, and has a diameter of 30 cm to 100 cm and a width of 15 cm to 100 cm, for example. The cooling roll 7 can be rotated at a predetermined rotational speed by a driving device (not shown). By controlling this rotational speed, the peripheral speed of the cooling roll 7 can be arbitrarily adjusted. The cooling speed by the strip casting apparatus is about 10 by selecting the rotation speed of the cooling roll 7 or the like. ² ° C / second to about 10 ^Five It can be controlled in the range of ° C / second.
[0091]
The surface of the chute 5 for guiding the molten metal is inclined at an angle (inclination angle) α with respect to the horizontal direction, and the distance between the tip of the chute 5 and the surface of the cooling roll is kept at several mm or less. And the chute | shoot 5 is arrange | positioned so that the line which connects the front-end | tip part and the center of the cooling roll 7 may form angle (beta) (0 degrees <= beta <= 90 degrees) with respect to a horizontal direction. The inclination angle α of the chute 5 is preferably 1 ° ≦ α ≦ 80 °, and more preferably satisfies the relationship of 5 ° ≦ α ≦ 60 °. The angle β preferably satisfies the relationship of 10 ° ≦ β ≦ 55 °.
[0092]
The molten metal 3 supplied onto the chute 5 is supplied from the tip of the chute 5 to the surface of the cooling roll 7 to form a molten metal paddle 6 on the surface of the cooling roll 7.
[0093]
The chute 5 can rectify the flow of the molten metal 3 by temporarily storing the molten metal 3 continuously supplied at a predetermined flow rate from the melting furnace 1 so as to delay the flow rate. If a blocking plate capable of selectively blocking the flow of the molten metal surface portion of the molten metal 3 supplied to the chute 5 is provided, the rectifying effect can be further improved. By using the chute 5, the molten metal 3 can be supplied in a state where the cooling roll 7 is widened to a substantially uniform thickness over a certain width in the trunk length direction (axial direction: perpendicular to the paper surface). By adjusting the inclination angle α of the molten metal guide surface of the chute 5, the molten metal supply speed can be finely adjusted. The molten metal flows on the inclined guide surface of the chute 5 by its own weight, and has a momentum component parallel to the horizontal direction (X-axis direction). As the inclination angle α of the chute 5 increases, the flow rate of the molten metal increases and the momentum also increases.
[0094]
In addition to the above function, the chute 5 also has a function of adjusting the temperature of the molten metal 3 immediately before reaching the cooling roll 7. The temperature of the molten metal 3 on the chute 5 is desirably a temperature that is 100 ° C. or more higher than the liquidus temperature. If the temperature of the molten metal 3 is too low, TiB adversely affects the alloy properties after quenching. ₂ This is because primary crystals such as nuclei are locally nucleated and may remain after solidification. On the other hand, if the molten metal temperature is too low, the viscosity of the molten metal increases and splash is likely to occur. The molten metal temperature on the chute 5 can be controlled by adjusting the molten metal temperature at the time of pouring from the melting furnace 1 into the chute 5, the heat capacity of the chute 5 itself, etc. 1 (not shown) may be provided.
[0095]
The chute 5 according to the present embodiment has a plurality of discharge portions provided at predetermined intervals along the axial direction of the cooling roll at the end arranged to face the outer peripheral surface of the cooling roll 7. ing. The width of the discharge part (width of one flow of the molten metal) is preferably set to 0.5 cm to 10.0 cm, and more preferably 0.7 cm to 4.0 cm. In this embodiment, the width of each molten metal flow in the discharge part is set to 1 cm. In addition, the width of the flow of the molten metal tends to spread in the lateral direction as the distance from the discharge portion increases. When a plurality of discharge portions are provided in the chute 5 to form a plurality of molten metal flows, it is preferable that adjacent molten metal flows are not in contact with each other.
[0096]
The molten metal 3 supplied onto the chute 5 is in contact with the cooling roll 7 along the axial direction of the cooling roll 7 and has substantially the same width as each discharge portion. Thereafter, the molten metal 3 that has come into contact with the cooling roll 7 with a predetermined discharge width moves on the circumferential surface of the roll as the cooling roll 7 rotates (so as to be pulled up by the cooling roll 7), and is cooled in this moving process. The In order to prevent molten metal leakage, the distance between the tip of the chute 5 and the cooling roll 7 is preferably set to 3 mm or less (particularly in the range of 0.4 to 0.7 mm).
[0097]
The gap between adjacent discharge portions is preferably set to 1 cm to 10 cm. Thus, if the molten metal contact part (molten cooling part) in the outer peripheral surface of the cooling roll 7 is isolate | separated into a several location, the molten metal discharged | emitted from each discharge part can be cooled effectively. As a result, a desired cooling rate can be achieved even when the amount of molten metal supplied to the chute 5 is increased.
[0098]
In addition, the form of the chute 5 is not limited to the above-described form, and may have a single discharge part, or the tapping width may be set larger.
[0099]
The molten alloy 3 solidified on the outer peripheral surface of the rotating cooling roll 7 becomes a strip-shaped solidified alloy 8 and peels from the cooling roll 7. In the case of this embodiment, the molten metal flowing out from each of the plurality of discharge portions becomes a band having a predetermined width and solidifies. The peeled solidified alloy 8 is crushed and collected by a collection device (not shown).
[0100]
In this way, the strip casting method does not use a nozzle like the melt spinning method, and there is no problem of injection speed limitation due to the nozzle diameter and clogging of the molten metal due to solidification at the nozzle part, so it is suitable for mass production. Yes. Moreover, since the heating equipment of the nozzle part and the pressure control mechanism for controlling the molten metal head pressure are not required, the initial equipment investment and the running cost can be reduced.
[0101]
In the melt spinning method, the nozzle part cannot be reused, so the nozzle with high processing cost had to be made disposable, but in the strip casting method, the chute can be used repeatedly, so the running cost Is cheap.
[0102]
Furthermore, according to the strip casting method, the cooling roll can be rotated at a slower speed than the melt spinning method, and the amount of the molten alloy hot water can be increased, so that the quenched alloy ribbon can be thickened.
[0103]
However, in the strip casting method, the molten alloy is not strongly injected onto the surface of the cooling roll. Therefore, when the cooling roll 7 rotates at a relatively high peripheral speed of 10 m / second or more, the molten metal paddle is formed on the surface of the cooling roll 7. There is a problem that it is difficult to form 6 stably. Further, when the nozzle is not used, since the pressure with which the molten alloy presses the roll surface is small, a minute gap is likely to occur between the molten alloy and the roll surface at the contact portion between the molten alloy and the roll surface. For this reason, the adhesiveness between the molten alloy and the roll surface is inferior in the strip casting method compared to the melt spinning method. The problem regarding the adhesion and the solution thereof will be described later.
[0104]
In the present embodiment, the upper limit value of the molten metal supply rate (processing amount) is defined by the supply rate per unit contact width between the molten metal and the cooling roll. In the case of the strip casting method, since the molten metal comes into contact with the cooling roll so as to have a predetermined contact width along the axial direction of the cooling roll, the cooling condition of the molten metal greatly depends on the molten metal supply speed per unit contact width.
[0105]
If the molten metal supply rate is too high, the cooling rate of the molten metal by the cooling roll decreases, and as a result, a rapidly cooled alloy containing a large amount of crystallized structure is produced without promoting amorphization, which is suitable for a nanocomposite magnet. It becomes impossible to obtain a raw material alloy. For this reason, in this invention, the supply rate (kg / min) per unit contact width (cm) is set to 3 kg / min / cm or less.
[0106]
As described above, for example, when the molten metal is brought into contact with the cooling roll in a contact form having a contact width of about 2 cm × 3, by setting the supply speed to about 0.5 kg / min / cm or more, about 3 kg / min The above processing amount can be realized.
[0107]
Thus, by supplying the molten metal at the supply speed in the specific range to the cooling roll rotating at the peripheral speed in the specific range, a desired quenched alloy can be produced with high productivity even when the strip cast method is used. can do. The strip casting method does not use a nozzle that significantly increases the manufacturing cost unlike the jet casting method, so the cost for the nozzle is unnecessary, and the production does not stop due to a nozzle clogging accident.
[0108]
In the present embodiment, the peripheral speed of the cooling roll can be set to 5 m / second or more and less than 20 m / second. If the roll peripheral speed is less than 5 m / sec, a desired quenched alloy cannot be obtained due to insufficient cooling capacity, and if it is 20 m / sec or more, it becomes difficult to pull up the molten metal by the roll, and the cooling alloy is in the form of flakes It may be difficult to collect because it is scattered. The optimum peripheral speed may vary depending on the structure, material, molten metal supply speed, etc. of the cooling roll, but if the peripheral speed is high, the resulting ribbon-like alloy becomes extremely thin and bulky, which makes it difficult to handle. On the other hand, when the peripheral speed is too high, the shape of the magnetic powder produced by pulverizing the ribbon-like alloy becomes flat, and therefore, the flowability of the magnetic powder and the cavity filling rate decrease when the magnetic powder is formed. As a result, the magnetic powder density of the magnet decreases, and the magnet characteristics deteriorate. On the other hand, when the peripheral speed is low, it is difficult to obtain a sufficient cooling speed. Therefore, the peripheral speed of the cooling roll is preferably set to 5 m / second or more and 20 m / second or less, more preferably 6 m / second or more and 15 m / second or less. A more preferable range of the peripheral speed of the cooling roll is 10 m / sec or more and 13 m / sec or less.
[0109]
If the supply rate per unit contact width exceeds 3 kg / min / cm, a predetermined cooling rate cannot be obtained, and it becomes difficult to produce a desired quenched alloy. The appropriate range of the supply speed per unit contact width may vary depending on the roll peripheral speed, roll structure, etc., but is preferably 2 kg / min / cm or less, and is 1.5 kg / min / cm or less. Is more preferable.
[0110]
Further, the molten metal supply speed (processing speed) of the entire apparatus is set to 3 kg / min or more because productivity is poor and inexpensive raw material supply cannot be realized if it is less than 3 kg / min. For this purpose, it is preferable that the supply rate per unit contact width is 0.4 kg / min / cm or more when the shape of the chute or the cooling roll is appropriately selected.
[0111]
For example, when a Cu roll having a diameter of about 35 cm and a width of about 15 cm is used, if the roll peripheral speed is 5 m / second to 10 m / second, the supply speed per unit contact width is 0.5 kg / minute / cm to 2 kg. About / min / cm is preferable. In this case, the rapid cooling step can be performed at a supply rate of 0.5 kg / min to 6 kg / min.
[0112]
By appropriately selecting the shape of the chute 5, the width and number of the molten metal discharge section, the molten metal supply speed, etc., the thickness (average value) and width of the obtained ribbon-like quenched alloy can be adjusted within an appropriate range. The width of the ribbon-like quenched alloy is preferably in the range of 15 mm to 80 mm. Moreover, if the thickness of the ribbon-like alloy is too thin, the bulk density becomes low, which makes it difficult to recover. If it is too thick, the cooling rate differs between the roll contact surface of the molten metal and the free surface (molten surface). Is not preferable because sufficient cannot be obtained. For this reason, the thickness of the ribbon-shaped alloy is preferably 50 μm or more and 250 μm or less, and more preferably 60 μm or more and 200 μm or less. A more preferable range of the thickness of the quenched alloy is 70 μm or more and 90 μm or less. In consideration of the packing density of the bonded magnet, the thickness of the quenched alloy preferably exceeds 80 μm.
[0113]
[Heat treatment]
In the present embodiment, the heat treatment is performed in an argon atmosphere. Preferably, the heating rate is 5 ° C./second to 20 ° C./second, and the temperature is maintained at 550 ° C. to 850 ° C. for 30 seconds to 20 minutes, and then cooled to room temperature. By this heat treatment, fine crystals of metastable phase precipitate and grow in the remaining amorphous phase, and a nanocomposite structure is formed. According to the present invention, R which is already fine at the time (as-cast) before the start of heat treatment. ₂ Fe ₁₄ B crystal phase (Nd ₂ Fe ₁₄ (B-type crystal phase) is present in an amount of 60% by volume or more of the whole, so that the coarsening of the α-Fe phase and other crystal phases is suppressed, and Nd ₂ Fe ₁₄ Each constituent phase (soft magnetic phase) other than the B-type crystal phase is uniformly refined. R after heat treatment ₂ Fe ₁₄ B crystal phase (Nd ₂ Fe ₁₄ The volume ratio of the B-type crystal phase) in the alloy is 65 to 85%.
[0114]
When the heat treatment temperature is lower than 550 ° C., a large amount of amorphous phase remains even after the heat treatment, and the coercive force may not reach a sufficient level depending on the rapid cooling conditions. Moreover, when the heat treatment temperature exceeds 850 ° C., the grain growth of each constituent phase is remarkable, and the residual magnetic flux density B _r Decreases, and the squareness of the demagnetization curve deteriorates. For this reason, the heat treatment temperature is preferably 550 ° C. or more and 850 ° C. or less, but the more preferable heat treatment temperature range is 570 ° C. or more and 820 ° C. or less.
[0115]
In the present invention, a sufficient amount of Nd in the quenched alloy. ₂ Fe ₁₄ The B-type compound phase is precipitated uniformly and finely. For this reason, even when the crystallization heat treatment is not performed on the quenched alloy, the rapidly solidified alloy itself can exhibit sufficient magnet characteristics. Therefore, the crystallization heat treatment is not an essential step in the present invention, but it is preferable to perform this for improving the magnet characteristics. It should be noted that the magnet characteristics can be sufficiently improved even by heat treatment at a lower temperature than in the past.
[0116]
In order to prevent oxidation of the alloy, the heat treatment atmosphere is Ar gas of 50 kPa or less or N ₂ An inert gas such as a gas is preferred. The heat treatment may be performed in a vacuum of 0.1 kPa or less.
[0117]
In the quenched alloy before heat treatment, R ₂ Fe ₁₄ In addition to the B compound phase and the amorphous phase, Fe _Three B phase, Fe _{twenty three} B ₆ , R ₂ Fe ₁₄ Phase B and R ₂ Fe _{twenty three} B _Three Metastable phases such as phases may be included. In that case, the R ₂ Fe _{twenty three} B _Three The phase disappears and R ₂ Fe ₁₄ Iron-based borides exhibiting a saturation magnetization equivalent to or higher than the saturation magnetization of the B phase (for example, Fe _{twenty three} B ₆ ) And α-Fe can be grown. In this specification, “Fe _Three "Phase B" means "Fe _3.5 Including “Phase B”.
[0118]
In the case of the present invention, even if a soft magnetic phase such as α-Fe is finally present, since the average crystal grain size of the soft magnetic phase is smaller than the average crystal grain size of the hard magnetic phase, Since the hard magnetic phase is magnetically coupled by exchange interaction, excellent magnetic properties are exhibited.
[0119]
Nd after heat treatment ₂ Fe ₁₄ The average crystal grain size of the B-type compound phase needs to be uniaxial crystal grain size of 300 nm or less, preferably 20 nm or more and 200 nm or less, and more preferably 20 nm or more and 150 nm or less. On the other hand, when the average crystal grain size of the ferromagnetic iron-based boride phase or α-Fe phase exceeds 50 nm, the exchange interaction acting between each constituent phase is weakened, and the squareness of the demagnetization curve is deteriorated. (BH) _max Will fall. Usually, these phases do not become precipitates having a diameter smaller than 1 nm, but precipitates having a size of several nm. From the above, the average crystal grain size of the soft magnetic phase such as boride phase or α-Fe phase is preferably 1 nm or more and 50 nm or less, and more preferably 5 nm or more and 30 nm or less. Nd due to magnetic properties ₂ Fe ₁₄ More preferably, the average crystal grain size of the B-type compound phase is 20 nm to 100 nm, and the average crystal grain size of the soft magnetic phase is 1 nm to 30 nm. In addition, in order to exhibit excellent performance as a replacement spring magnet, Nd ₂ Fe ₁₄ The average crystal grain size of the B-type compound phase is preferably larger than the average crystal grain size of the soft magnetic phase.
[0120]
Further, according to the present embodiment, as shown in FIG. ₂ Fe ₁₄ A structure in which a fine iron-based boride phase ((Fe, Ti) -B compound) is present at the grain boundary or subgrain boundary of the B-type compound phase is obtained. Such a structure is suitable for maximizing exchange interaction between constituent phases. Ti is present in the iron-based boride. This is considered to be because Ti has a strong affinity for B and Ti is easily concentrated in the iron-based boride. Since Ti and B are strongly bonded in the iron-based boride, the addition of Ti is considered to stabilize the iron-based boride.
[0121]
Note that the quenched alloy ribbon may be roughly cut or pulverized before the heat treatment. If the obtained magnet is pulverized after heat treatment to produce magnet powder (magnetic powder), various bonded magnets can be produced from the magnetic powder by known processes. When producing a bonded magnet, the iron-based rare earth alloy magnetic powder is mixed with an epoxy resin or a nylon resin and formed into a desired shape. At this time, the nanocomposite magnetic powder may be mixed with other types of magnetic powder, for example, Sm—Fe—N magnetic powder or hard ferrite magnetic powder.
[0122]
Various rotary machines, such as a motor and an actuator, can be manufactured using the above-mentioned bond magnet.
[0123]
When the magnetic powder obtained by the method of the present invention is used for an injection-molded bonded magnet, it is preferably pulverized so that the average particle size is 200 μm or less, and more preferably the average particle size of the powder is 30 μm or more and 150 μm or less. . Moreover, when using for a compression-molded bond magnet, it is preferable to grind | pulverize so that a particle size may be 300 micrometers or less, and the average particle diameter of a more preferable powder is 30 micrometers or more and 250 micrometers or less. A more preferable range is 50 μm or more and 200 μm or less.
[0124]
(Embodiment 2)
Next, a second embodiment of the present invention will be described.
[0125]
When a molten alloy of the above composition containing Ti as an essential element is quenched and solidified by a strip cast method, a compound in which Ti and B are combined (TiB ₂ Etc.) is easily formed in the molten metal, and as a result, the liquidus temperature of the molten metal is higher than that of the molten iron-based rare earth magnet material alloy having the conventional composition. If the liquidus temperature of the molten metal increases, the molten metal temperature is set higher (for example, about 100 ° C. higher than the liquidus temperature), and if the melt viscosity is not kept sufficiently low, It cannot be realized.
[0126]
However, when the molten alloy is rapidly cooled and solidified on the surface of the cooling roll, if the hot water temperature is raised, the roll surface temperature will rise, so the ribbon of the quenched alloy will not easily peel off from the cooling roll, and it will be easy to wind around the cooling roll. Become. When the alloy ribbon is wound around the roll, the molten metal is successively supplied onto the wound alloy, and the crystal phase generated in the quenched alloy becomes coarse, resulting in deterioration of the final magnet characteristics. Become.
[0127]
This problem hardly occurs in the melt spinning method in which a relatively small amount of molten alloy is injected from a nozzle. In the case of the melt spinning method, the amount of the molten metal that contacts the surface of the cooling roll is small, and the adhesion between the molten metal that is strongly jetted and the roll surface is good. As a result, the ability of the roll to cool the molten metal is unlikely to decrease, and cooling of the molten metal proceeds uniformly and sufficiently.
[0128]
On the other hand, in the case of the strip casting method, since no nozzle is used, it is difficult to uniformly and sufficiently cool a large amount of molten alloy. Further, in the alloy composition used in the present invention, the cooling rate of the molten metal and the cooling uniformity greatly influence the microstructure of the quenched alloy, and determine the magnet characteristics. For this reason, in order to mass-produce high-performance nanocomposite magnets by the strip cast method, it is necessary to sufficiently prevent winding of the alloy ribbon around the cooling roll.
[0129]
The present inventor has found that the liquidus temperature of the molten alloy is lowered by 10 ° C. or more (for example, about 40 to 80 ° C.) by adding an appropriate amount of Nb to the alloy having the composition system described above. When the liquidus temperature of the molten alloy is lowered, even if the molten metal temperature is lowered by that amount, the viscosity of the molten metal is hardly increased, and stable tapping can be continuously performed. When the temperature of the hot water is lowered, sufficient cooling can be achieved on the surface of the cooling roll, so that it is possible to prevent winding of the roll and to make the rapidly solidified alloy structure uniform and fine.
[0130]
Therefore, in this embodiment, the composition formula is (Fe _1-m T _m ) _100-xyzn (B _1-p C _p ) _x R _y Ti _z Nb _n The alloy melt expressed by is quenched by strip casting. Here, T is expressed by one or more elements selected from the group consisting of Co and Ni, R is expressed by one or more elements selected from the group consisting of Y (yttrium) and rare earth metals), and the composition ratio x , Y, z, m, n, and p each satisfy the following relational expression.
[0131]
10 <x ≦ 25 atomic%
7 ≦ y <10 atomic%
0.5 ≦ z ≦ 12 atomic%
0 ≦ m ≦ 0.5
0.1 ≦ n ≦ 5 atomic%, and
0 ≦ p ≦ 0.25
[0132]
In order to prevent the alloy from being wound by the cooling roll, it is preferable not only to add Nb but also to adjust the atmospheric gas pressure to an appropriate range as described above.
[0133]
In this embodiment, a rapidly solidified alloy is produced using the strip casting apparatus shown in FIG. In order to prevent oxidation of the raw material alloy containing rare earth elements R and Fe which are easily oxidized, the alloy manufacturing process is performed in an inert gas atmosphere. As the inert gas, a rare gas such as helium or argon or nitrogen can be used.
[0134]
The strip casting apparatus of FIG. 3 is disposed in a chamber that can be evacuated in an inert gas atmosphere. As in the apparatus of FIG. 1, this strip casting apparatus includes a melting furnace 1 for melting an alloy raw material, a cooling roll 7 for rapidly cooling and solidifying a molten alloy 3 supplied from the melting furnace 1, and a melting furnace A chute (tundish) 5 that guides the molten metal 3 from 1 to the cooling roll 7 and a scraper gas ejector 9 that solidifies and easily peels the ribbon-shaped alloy 8 from the cooling roll 7 are provided.
[0135]
The melting furnace 1 can supply the molten metal 3 produced by melting the alloy raw material to the chute 5 at a substantially constant supply amount. This supply amount can be arbitrarily adjusted by controlling the operation of tilting the melting furnace 1.
[0136]
The outer peripheral surface of the cooling roll 7 is formed of a material having good thermal conductivity such as copper, and has a diameter (2r) of 30 cm to 100 cm and a width of 15 cm to 100 cm. The cooling roll 7 can be rotated at a predetermined rotational speed by a driving device (not shown). By controlling this rotational speed, the peripheral speed of the cooling roll 7 can be arbitrarily adjusted. The cooling speed by the strip casting apparatus is about 10 by selecting the rotation speed of the cooling roll 7 or the like. ² ° C / second to about 2 × 10 ^Four It can be controlled in the range of ° C / second.
[0137]
The molten metal 3 supplied onto the chute 5 is supplied from the tip of the chute without applying pressure to the surface of the cooling roll 7, and a molten metal paddle 6 is formed on the surface of the cooling roll 7.
[0138]
The chute 5 is made of ceramics or the like, can delay the flow rate so as to temporarily store the molten metal 3 continuously supplied from the melting furnace 1 at a predetermined flow rate, and can rectify the flow of the molten metal 3. . If a blocking plate capable of selectively blocking the flow of the molten metal surface portion of the molten metal 3 supplied to the chute 5 is provided, the rectifying effect can be further improved.
[0139]
Various conditions in the strip casting process apply as described for the first embodiment. Further, the subsequent process performed on the quenched alloy is the same as the process in the first embodiment.
[0140]
According to the present embodiment, by adding Nb together with Ti to the iron-based rare earth alloy, the liquidus temperature of the molten alloy can be lowered, and the rapidly cooled alloy can be stably produced at the mass production level.
[0141]
Note that the composition ratio of Nb is preferably 0.1 atomic percent or more and 5 atomic percent or less, and more preferably 0.5 atomic percent or more and 3 atomic percent or less.
[0142]
(Embodiment 3)
Next, a third embodiment of the present invention will be described.
[0143]
As described above, in the alloy used in the present invention, a compound in which Ti and B are bonded (TiB ₂ Etc.) is easily formed in the molten metal, and as a result, the liquidus temperature of the molten metal is higher than that of the molten iron-based rare earth magnet material alloy having the conventional composition.
[0144]
The present inventor has found that if an appropriate amount of C (carbon) is added to an iron-based alloy containing Ti and B, the liquidus temperature of the molten alloy decreases by 5 ° C. or more (for example, about 10 to 40 ° C.). . When the liquidus temperature of the molten alloy decreases due to the addition of carbon, even if the molten metal temperature is lowered accordingly, TiB ₂ Thus, the melt viscosity hardly increases and a stable melt flow can be continuously formed. When the molten metal temperature is lowered, sufficient cooling can be achieved on the surface of the cooling roll, so that the roll can be prevented from being wound and the rapidly solidified alloy structure can be uniformly refined.
[0145]
In this embodiment, the composition formula is (Fe _1-m T _m ) _100-xyzn (B _1-p C _p ) _x R _y Ti _z M _n An iron-based rare earth rapidly solidified alloy is produced by quenching the molten alloy expressed by Here, T is one or more elements selected from the group consisting of Co and Ni, and R is one or more elements selected from the group consisting of Y (yttrium) and rare earth metals. M is one or more selected from the group consisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au, and Pb It is an element.
[0146]
Each of x, y, z, m, n, and p in the composition formula satisfies the following relational expression.
[0147]
10 <x ≦ 25 atomic%
7 ≦ y <10 atomic%
0.5 ≦ z ≦ 12 atomic%
0 ≦ m ≦ 0.5
0 ≦ n ≦ 10 atomic%, and
0.01 ≦ p ≦ 0.25
[0148]
In order to solidify the molten metal of the alloy, the strip casting apparatus shown in FIG. 3 is also used in this embodiment. In the present embodiment, a raw material having an oxygen concentration of 1000 ppm or less by mass ratio is melted, and the oxygen concentration of the alloy in the molten state is controlled to 3000 ppm or less by mass ratio. Since the oxygen concentration of the molten metal changes depending on the oxygen partial pressure in the atmosphere and the time from melting to rapid solidification, in this embodiment, by adjusting these conditions, the oxygen concentration does not exceed 3000 ppm. Yes.
[0149]
The molten metal 3 supplied onto the chute 5 is supplied without applying pressure to the surface of the cooling roll 7 from the tip of the chute, and forms a molten paddle 6 on the surface of the cooling roll 7. In this embodiment, since the liquidus temperature of the molten metal is kept low by adding carbon, the kinematic viscosity of the molten metal is 5 × 10 5 when the molten metal temperature is 1200 ° C. or higher. ^-6 m ² / Second or less, and a smooth hot water flow is realized.
[0150]
The temperature of the molten metal 3 on the chute 5 is desirably a temperature that is 100 ° C. or more higher than the liquidus temperature. This is because if the temperature of the molten metal 3 is too low, primary crystals that adversely affect the properties of the alloy after quenching are locally nucleated and may remain after solidification.
[0151]
Various conditions in the strip casting process apply as described for the first embodiment. Further, the subsequent process performed on the quenched alloy is the same as the process in the first embodiment.
[0152]
In addition, since the “bulk density” of the obtained quenched alloy ribbon is often 0.5 g (g) / cc or less, after quenching, the “bulk density” is 1 g / cc using an appropriate pulverizer. It is preferable to grind and collect the alloy so that it is as described above.
[0153]
The steps after the strip casting step may be the same as the steps in the first embodiment.
[0154]
According to the present embodiment, by adding C together with Ti to the iron-based rare earth alloy, the liquidus temperature of the molten alloy can be lowered, and the rapidly cooled alloy can be stably produced at the mass production level.
[0155]
(Embodiment 4)
The peripheral speed of the cooling roll in the conventional strip casting method is very slow, about 1-2 m / sec. In the alloy composition used in the present invention, a high-quality quenched alloy structure can be formed at a relatively low speed by adding Ti. However, in order to improve the magnet characteristics as much as possible, it is much higher than the peripheral speed in the conventional strip casting method. It is preferable to cool the molten metal at a high peripheral speed.
[0156]
However, in the strip casting method, when the rotation speed of the cooling roll is increased, it becomes difficult to sufficiently raise the molten alloy. In the case of the strip casting method, the adhesion of the molten metal to the rotating roll surface is lower than that in the melt spinning method. This is partly because a thin air layer formed on the roll surface penetrates between the molten metal and the roll surface. For this reason, when the cooling roll is rotated at a high speed, the molten metal slides on the roll surface, and the molten metal cannot be pulled up. On the other hand, in the case of the melt spinning method, a thin jet of molten metal having a large momentum hits the cooling roll surface through the nozzle orifice, so that the molten metal can be brought into close contact with the roll surface by breaking the air layer. Even when the cooling roll rotates at a high speed, it is possible to form a desired rapidly solidified alloy.
[0157]
Under these circumstances, conventionally, when it is necessary to increase the cooling rate, the peripheral speed of the cooling roll has been set to be higher (for example, 20 m / second or more) using the melt spinning method. On the contrary, when the cooling rate may be slow, the peripheral speed of the cooling roll is set low (for example, 1 to 2 m / sec) using the strip casting method.
[0158]
When an iron-based rare earth alloy magnet is produced by a liquid quenching method, a desired microstructure cannot be obtained unless the cooling rate is sufficiently increased. In particular, when producing a nanocomposite magnet in which a hard magnetic phase composed of an Nd—Fe—B compound and a soft magnetic phase such as α-Fe are magnetically coupled by an exchange interaction, conventionally, the melt spinning method must be used. In this case, the desired cooling structure could not be obtained because the cooling rate was insufficient. Therefore, mass production of such a nanocomposite magnet by the strip cast method has not been realized.
[0159]
In order to achieve a cooling speed that has been considered difficult to achieve by the conventional strip casting method, the present inventors have used a strip casting method using a cooling roll that rotates at a high speed (peripheral speed: 10 m / second or more). Development has been considered. In the strip casting method of the present inventors, the molten alloy is supplied onto an inclined chute (guide means), and a lateral flow of the molten alloy is formed on the chute by utilizing the weight of the molten metal. By imparting a relatively large momentum to the molten metal in this way, it is possible to hit the molten metal against the surface of the cooling roll and to bring the molten metal into close contact with the surface of the cooling roll rotating at high speed.
[0160]
However, according to the above-described strip casting method of the present inventors, it has been found that even when C or Nb is added and the molten metal is rapidly cooled in a reduced pressure atmosphere, the quenched alloy may be wound around the cooling roll. . When the quenching alloy is wound around the cooling roll, the quenching process is forced to be interrupted, and the quenching process cannot be continued. This is a big hindrance to the realization of mass production.
[0161]
In the present embodiment, the structure of the chute and the cooling roll, which are useful for stably forming the molten metal paddle on the cooling roll rotating at high speed and preventing the quenching alloy from being wound around the roll, will be described in detail. To do.
[0162]
In the present embodiment, the strip casting method is performed using the apparatus shown in FIG. As described above, the molten metal guide surface of the chute 5 is inclined with respect to the horizontal direction to form a molten metal flow path to the cooling roll 7. The angle (inclination angle) α between the guide surface of the chute 5 and the horizontal direction is an important parameter for finely controlling the supply amount (rate) of the molten metal.
[0163]
The molten metal 3 supplied onto the chute 5 is supplied from the tip of the chute 5 with a momentum in the horizontal direction with respect to the surface of the cooling roll 7, and forms a melt paddle 6 on the surface of the cooling roll 7.
[0164]
FIG. 4 is a perspective view showing the upper surface of the chute 5. The chute 5 has a guide for guiding the molten metal received at one place to the tip. Some of these guides exist not only on both sides of the flow path but also in the central portion, and the molten metal flow can be divided into two strips. In the example of FIG. 4, the width of each of the two molten metal flows is defined as 10 mm, and the molten metal of each filament is supplied to the surface of the cooling roll at an interval of 10 mm. According to the chute 5 having such a guide, the molten metal 3 is spread in a substantially uniform thickness over a certain width in the body length direction of the cooling roll 7 (axial direction: direction perpendicular to the paper surface of FIG. 3). Can be supplied. In addition, the width | variety (width | variety of a quenching alloy) of each ribbon at this time is set to 5-20 mm. This is because mass productivity decreases when the ribbon width is less than 5 mm, and stable casting becomes difficult when the ribbon width exceeds 20 mm.
[0165]
Usually, fine irregularities exist on the surface of the cooling roll 7 used in the strip casting apparatus. When the surface roughness of the chill roll 7 is increased, the effective contact area between the molten alloy and the surface of the chill roll 7 is reduced due to the fine recesses present on the roll surface.
[0166]
FIG. 5 schematically shows the cross-sectional shape of the molten metal that contacts the surface of the cooling roll 7 that rotates at a peripheral speed of 10 m / sec. An atmosphere gas is caught between the surface of the cooling roll 7 and the lower surface of the molten metal, and a large number of air pockets 50 are formed. The larger the surface roughness of the cooling roll 7, the lower the effective contact area between the roll surface and the molten metal. As a result, the amount of heat removed from the molten metal by the cooling roll 7 is reduced, and the cooling rate of the molten alloy 3 is substantially reduced. When the cooling capacity of the cooling roll 7 is reduced in this way, the temperature of the quenched alloy 8 that is solidifying in contact with the cooling roll 7 is not sufficiently lowered.
[0167]
The rapidly cooled alloy 8 shrinks during solidification, but as the solidification shrinkage becomes insufficient, the rapid cooling alloy 8 is less likely to be peeled off from the rotating cooling roll 7 and is easily wound around the cooling roll 7. When the ribbon-like quenching alloy 8 is wound around the cooling roll 7, the cooling process cannot be continued. In particular, in the case of the strip casting method, there is a problem that the rapidly cooled alloy 8 is likely to be wound around the cooling roll 7 because the roll circumferential size of the portion where the molten metal is in contact with the roll surface is longer than that in the melt spinning method.
[0168]
On the other hand, according to the melt spinning method, as shown in FIG. 6, a relatively small amount of molten alloy is sprayed onto the surface of the cooling roll 7 through a nozzle, and the molten metal is pressed against the roll surface. Even when the degree is high, the adhesion between the roll surface and the molten metal is good and the cooling capacity is high, so that it is easy to uniformly cool the molten alloy at a sufficient speed.
[0169]
As described above, when the strip casting method of the type employed in the present invention is used, when the cooling roll 7 is rotated at a high speed at a peripheral speed of 10 m / second or more, the center line roughness Ra on the surface of the cooling roll 7 is rapidly quenched. 8 has an important influence on the winding around the cooling roll 7. According to the experiment of the present inventor, if the center line roughness Ra on the surface of the cooling roll 7 is 20 μm or less, a sufficient cooling effect can be obtained, so that the quenched alloy 8 can be prevented from being wound around the cooling roll 7. I understood.
[0170]
From the above, in the present invention, the center line roughness Ra of the cooling roll surface is set to 20 μm or less. In order to continue stable operation at the mass production level, the center line roughness Ra of the cooling roll surface is preferably set to 13 μm or less, and more preferably set to 7 μm or less.
[0171]
Further, in the strip casting method employed in the present invention, as shown in FIGS. 3 and 4, the molten alloy 3 flows slowly on the inclined chute 5, so that an appropriate paddle is formed on the surface of the cooling roll 7 that rotates at a high speed. In order to form 6, the magnitude of the kinematic viscosity of the molten alloy 3 is important. According to experiments, the kinematic viscosity of the molten alloy 3 is 5 × 10. ^-6 m ² It has been found that when the speed exceeds / sec, the paddle 6 is not formed on the cooling roll 7, and the molten metal 3 becomes splash and cannot be rapidly cooled. For this reason, the kinematic viscosity of the molten alloy 3 is 5 × 10. ^-6 m ² Is preferably adjusted to 1 / second or less. ^-7 m ² It is more preferable to adjust to less than / sec.
[0172]
If the surface temperature of the chute 5 is too low, the kinematic viscosity of the molten alloy 3 may become too high before it reaches the cooling roll 7. When the surface temperature of the chute 5 is 300 ° C. or lower, the molten metal is cooled on the chute 5 and the kinematic viscosity is 5 × 10. ^-6 m ² Therefore, the surface temperature of the chute 5 is preferably kept at 300 ° C. or higher. The surface temperature of the chute 5 is preferably maintained at 450 ° C. or higher, more preferably 550 ° C. or higher.
[0173]
As the material of the chute 5, boron nitride (BN) can be used in addition to ceramic materials such as alumina, silica, zirconia, magnesia and mullite. Alumina (Al) which has excellent “wetability” with molten iron-based rare earth alloys and is difficult to react with rare earths ₂ O _Three ) Is preferably used. Moreover, in order to prevent the chute 5 from being broken by heat shock, porous ceramics are preferable to dense materials. However, the surface of the chute through which the molten metal flows is preferably as smooth as possible.
[0174]
In order to stably quench the molten alloy with the cooling roll 7, it is preferable to produce the cooling roll using a base material having a thermal conductivity of 50 W / m / K or more. As the base material of the cooling roll 7, iron, carbon steel, tungsten, molybdenum, beryllium, and tantalum can be used in addition to copper and a copper alloy. In order to stably cool the molten metal, it is particularly preferable to use copper and a copper alloy, tungsten, molybdenum, or beryllium having a thermal conductivity of 100 W / m / K or more.
[0175]
It is preferable to coat the surface of the base material of the cooling roll 7 with chromium, nickel having a thickness of 1 μm to 100 μm or a combination thereof. Thereby, the defect of a cooling roll base material with low melting points, such as copper, and low hardness can be compensated. Further, melting and scratching of the roll base material that occurs on the roll surface during molten metal cooling can be suppressed. As a result, the center line roughness Ra of the roll surface can be maintained at 20 μm or less for a long time. The thickness of the plating film is preferably in the range of 1 μm to 100 μm from the viewpoint of film strength and heat conduction. The more preferable thickness of the plating film is 5 μm to 70 μm, and the most preferable thickness is 10 μm to 40 μm.
[0176]
In addition, when the molten metal quenching process speed | rate per 1 piece of molten alloy 3 is less than 1 kg / min, the paddle 6 is not formed on a cooling roll, but the stable molten metal quenching state cannot be maintained. On the other hand, when the molten metal quenching speed per strip of the molten alloy 3 is 4 kg / min or more, the molten metal 3 is supplied to the volume of the paddle 6 that can be formed on the roll surface, so that the excess molten metal 3 is splashed and not rapidly cooled. . Therefore, it is preferable that the molten metal quenching treatment rate per strip of the molten alloy 3 is 0.7 kg / min or more and less than 4 kg / min. A more preferable range is 1 kg / min or more and less than 3 kg / min, and a most preferable range is 1 kg / min or more and less than 2 kg / min. From the viewpoint of mass productivity, it is preferable to use a guide as shown in FIG. When flowing a plurality of melts, it is desirable to provide an appropriate interval so that the melts do not contact each other.
[0177]
In this embodiment, the roll surface speed of the cooling roll 7 is set to 10 m / sec or more and 26 m / sec or less. By setting the roll surface speed to 10 m / sec or more, the α-Fe phase can more effectively suppress precipitation. However, when the roll surface peripheral speed exceeds 26 m / sec, the molten metal paddle 6 to be generated on the roll is not stabilized, and the molten metal jumps off (splash occurs). The molten metal cannot be quenched. A more preferable upper limit of the roll surface speed is 23 m / sec or less, and a more preferable upper limit is 20 m / sec or less.
[0178]
The generation state of the paddle 6 is influenced not only by the roll surface speed but also by the molten metal supply speed to the cooling roll 7. In order to maintain a stable generation state of the paddle 6, it is preferable to adjust the molten metal supply speed per one flow of the molten metal supplied to the cooling roll 7 within the above-described range.
[0179]
In the present embodiment, the pressure of the quenching atmosphere is adjusted to 0.13 kPa or more and less than 100 kPa. If the pressure in the quenching atmosphere is less than 0.13 kPa, the molten alloy may stick to the surface of the cooling roll, and the quenched alloy may not be peeled from the roll. On the other hand, when the pressure of the quenching atmosphere exceeds 100 kPa, the atmospheric gas is caught between the surface of the cooling roll and the molten alloy, and gas pockets are easily generated. When the gas pocket is formed, a uniform quenching state cannot be obtained, and a heterogeneous quenching structure is obtained, so that a supercooled state cannot be stably obtained. A preferable pressure range of the quenching atmosphere is 1.3 kPa or more and 90 kPa or less, a more preferable range is 10 kPa or more and 70 kPa or less, and a further preferable range is 10 kPa or more and 60 kPa or less. The most preferable range is 30 kPa or more and 50 kPa or less.
[0180]
When the molten alloy is rapidly cooled as described above, the adhesion of the molten alloy to the surface of the cooling roll is improved and a high cooling effect is uniformly imparted, so that the quenched alloy is appropriately formed and wound around the cooling roll. Trouble hardly occurs.
[0181]
[Structure structure of quenched alloy]
FIG. 7 schematically shows how the cross-sectional structure of the quenched alloy changes depending on whether or not Ti is added.
[0182]
First, as can be seen from FIG. 7, the quenched alloy (ribbon) produced by the strip casting method is thicker than the quenched alloy produced by the melt spinning method, so that the free surface of the quenched alloy (the surface not in contact with the cooling roll: Crystal grains are formed in the vicinity of the upper end surface. Moreover, crystal grains are also formed in the vicinity of the roll surface (surface in contact with the cooling roll: lower end surface). This is because heterogeneous nuclei are easily generated on the roll surface, and crystal growth easily proceeds around the heterogeneous nuclei. In the quenched alloy produced by the strip cast method, the size of crystal grains and the volume density of crystal grains become smaller as they approach the center of the film from each end face.
[0183]
When Ti is added, the formed crystal grains are generally small, and in particular, α-Fe is fine and few. An amorphous portion is likely to exist at the center of the film, and the crystalline layer formed on the roll surface side is thinner than the crystalline layer formed on the free surface side. Further, when Ti is added, iron-based boride (Fe-B) is precipitated. On the other hand, when Ti is not added, the size of the crystal grains is large, and α-Fe is particularly coarse. Since the cooling rate of the free surface decreases as the quenching alloy becomes thicker, coarse crystal grains are more likely to be formed on the free surface side as the quenching alloy becomes thicker. For this reason, the thicker the quenched alloy, the lower the finally obtained magnet properties. However, the addition of Ti has an effect of suppressing the coarsening of the crystal grains, so that it becomes easy to form a rapidly quenched alloy. In the case of this embodiment, it is possible to set the thickness of the quenched alloy to a range of about 50 to 200 μm. From the viewpoint of the shape of the powder particles after pulverization and the magnetic properties, the preferred thickness of the quenched alloy is 60 to 150 μm, and the more preferred thickness is 70 to 120 μm. As described above, according to the present invention, it is possible to produce a rapidly cooled alloy having a thickness of 80 μm or more, which was difficult in the prior art, and obtain a nanocomposite magnet having excellent magnetic properties. In FIG. 7, each crystal grain is schematically shown larger than the actual size. The actual size of each crystal grain is so small that it cannot be illustrated.
[0184]
According to the high-speed strip casting method according to the present embodiment, even if amorphous exists in the central portion of the cross-section of the quenched alloy, the crystalline portion is on the free surface and the roll surface (two end surfaces crossing the thickness direction). Exists. And when Ti is added, since the coarsening of (alpha) -Fe is suppressed, the magnet characteristic will be excellent. Since the peripheral speed of the cooling roll is much faster than the conventional method of strip casting, the crystal grains are not coarsened, and a rapidly cooled alloy having a structure suitable for a nanocomposite magnet can be obtained. Moreover, since the alloy (ribbon) after quenching has a structure in which crystal phases exist in the vicinity of both the free surface and the roll surface, even when the quenched alloy is pulverized before heat treatment, the quenched alloy is pulverized. Becomes easier and the pulverization efficiency is improved.
[0185]
When the nanocomposite magnet finally obtained by the manufacturing method of this embodiment is used for a motor, 600 kA / m is required to keep a sufficient level of magnetization even when a strong demagnetizing field acts on the magnet. High intrinsic coercive force H _cJ It is desirable to have In order to realize such a high coercive force, R contained in the microstructure of the quenched alloy. ₂ Fe ₁₄ It is necessary to make the volume ratio of the B-type compound phase 60% or more.
[0186]
The composition of the iron-based rare earth alloy in this embodiment is R ₂ Fe ₁₄ The R concentration is low and the B concentration is high compared to the stoichiometric composition of the B phase type compound. In such a composition, the addition of Ti facilitates the formation of iron-based borides by binding excess B to iron. The iron-based boride obtained by the addition of Ti has a size on the order of nanometers and is ferromagnetic. The addition of Ti not only suppresses the precipitation of coarse Fe, but also produces the fine ferromagnetic iron-based boride described above. ₂ Fe ₁₄ Without binding to the B-type compound phase by exchange interaction and causing a decrease in magnetization, R ₂ Fe ₁₄ It becomes possible to exhibit the same level of hard magnetic properties as an iron-based rare earth alloy magnet having the same stoichiometric composition as B.
[0187]
In the case of the iron-based rare earth alloy magnet in the present embodiment, R which is hard magnetic ₂ Fe ₁₄ In addition to the B phase, the saturation magnetization value is R ₂ Fe ₁₄ Because the same structure contains soft magnetic iron-based borides that are equal to or higher than the B phase, the magnet recoil permeability μ _r Is equivalent coercive force H _cJ In an alloy having the above, it is higher than that of an iron-based rare earth alloy magnet not containing an iron-based boride. Specifically, the recoil permeability μ in the iron-based rare earth alloy magnet of the present embodiment. _r Shows a value of 1.1 to 1.4 when the composition ratio y of the rare earth R is 8.5 atomic% or more and less than 10 atomic%, and the composition ratio y is 7 atomic% or more and 8.5 atomic% or less. Shows a value of 1.2 to 2.0. In the range where the composition ratio y is 8.5 atomic% or more and less than 10 atomic%, the residual magnetic flux density B of the magnet of the present embodiment. _r Is 0.7-0.9T, coercive force H _cJ Is 600 to 1200 kA / m, and in the range where the composition ratio y is 7 atomic% or more and 8.5 atomic% or less, the residual magnetic flux density B _r Is 0.75-0.95T, coercive force H _cJ Is 500 to 950 kA / m. Recoil permeability μ _r The measurement of. The method was described in JIS standard C2501-1989. Recoil permeability μ _r Is a parameter specific to an alloy in which a nanocomposite structure is formed, that is, an alloy in which a hard magnetic phase and a soft magnetic phase are crystallized and magnetically coupled by exchange interaction.
[0188]
Recoil permeability μ _r Is an important index for evaluating the performance of a magnet when the magnet is used in a motor. Hereinafter, this point will be described. That is, when the rotational speed of the motor is increased, the counter electromotive force is increased, and when the magnitude of the counter electromotive force becomes equal to the input voltage, the increase in the motor rotational speed is stopped. In order to further increase the rotational speed of the motor, it is necessary to lower the back electromotive force by electrically lowering the magnet operating point (-B / H) to the low permeance side (field weakening control). The effect of further increasing the upper limit of the motor rotational speed by such control is that the magnet recoil permeability μ _r The higher the value, the more prominent. The magnet according to the present invention has a high recoil permeability μ as described above. _r Therefore, it is preferably used for a motor.
[0189]
In the present invention, the peripheral speed of the cooling roll is much faster than that of the conventional strip casting method, but it is slower than the peripheral speed realized by the melt spinning method (for example, 20 m / second or more). If Ti is not added, α-Fe is R ₂ Fe ₁₄ Preferentially precipitates over the B compound and becomes coarse.
[0190]
The quenched alloy obtained by the quenching method described above is subjected to heat treatment after being pulverized.
[0191]
R in the alloy after heat treatment ₂ Fe ₁₄ The size (average particle size or average major axis length) of the B-type compound phase needs to be uniaxial crystal particle size of 300 nm or less, preferably 20 nm to 200 nm, and preferably 20 nm to 100 nm. More preferably. On the other hand, when the average crystal grain size of the iron-based boride phase or α-Fe phase exceeds 50 nm, the exchange interaction acting between the constituent phases is weakened, and the squareness of the demagnetization curve is deteriorated (BH). _max Will fall. When the average crystal grain size is less than 1 nm, a high coercive force cannot be obtained. From the above, the average particle size of the soft magnetic phase such as boride phase or α-Fe phase is preferably 1 nm or more and 50 nm or less, and more preferably 30 nm or less.
[0192]
The size (particle size) of the powder particles finally obtained from the alloy produced by the production method according to the present invention is preferably 10 to 300 μm, more preferably 50 to 150 μm. A more preferable range of the particle size is 80 to 110 μm.
[0193]
The average ratio (aspect ratio) of the minor axis size to the major axis size of the powder particles thus obtained is about 0.3 to 1.0. Since the thickness of the quenched alloy produced in the present embodiment is sufficiently thick with respect to the powder particle size, it is easy to obtain powder particles having a shape close to an equiaxed shape. On the other hand, since the thickness of the quenched alloy produced by ordinary melt spinning is as thin as about 20 to 40 μm, flaky powder particles having a small aspect ratio can be obtained under the same pulverization conditions as in this embodiment. Since the powder obtained in this embodiment has an aspect ratio close to 1, it has excellent filling properties and fluidity, and is optimal for bonded magnets.
[0194]
The coercive force H of the magnetic powder obtained in this way _cJ Can show a value of 600 kA / m or more.
[0195]
[Dependence of particle size distribution on oxidation resistance and magnetic properties of magnetic powder]
The oxidation resistance of the magnetic powder finally obtained from the alloy by the production method according to the present invention (hereinafter referred to as “nanocomposite magnetic powder”) and the particle size distribution dependence of the magnetic properties will be described in comparison with conventional quenched magnet powder.
[0196]
Here, the nanocomposite magnetic powder according to the present invention is compared with MQP-B and MQP-O (each having a maximum particle size of 300 μm or less) commercially available from MQI as conventional quenched magnet powder. A sample of nanocomposite magnetic powder according to the present invention was prepared as follows.
[0197]
First, a quenched alloy (Nd: 9 atomic%, B: 11 atomic%, Ti: 3 atomic%, Co: 2 atomic%, balance Fe alloy, average thickness: 70 μm, standard deviation σ: 13 μm) is pulverized to 850 μm or less, and a hoop belt furnace having a soaking zone with a length of about 500 mm is used and maintained at 680 ° C. at a belt feed rate of 100 mm / min under Ar flow. Heat treatment was performed by charging the powder at a feed rate of 20 g / min to obtain magnetic powder. The obtained magnetic powder was pulverized to a particle size distribution containing about 30% of a powder having an aspect ratio of 0.4 or more and 1.0 or less on a volume basis using a pin disk mill, and a sample NCP-0 of nanocomposite magnetic powder was obtained. did.
[0198]
Table 1 shows the oxygen content and magnetic properties after leaving each magnetic powder in the atmosphere at various temperatures (23 ° C., 300 ° C. and 350 ° C.) for 1 hour. Magnetic properties were measured using a vibrating magnetometer. Together with the results measured at 23 ° C., the results measured after being allowed to stand at 300 ° C. and 350 ° C. for 1 hour in the air are also shown.
[0199]
As shown in Table 1, MQP-B increased in oxygen content to 0.67 mass% when left in the atmosphere at 300 ° C. for 1 hour, and increased to 1.93 mass% when left at 350 ° C. for 1 hour. Increased. When MQP-O was left at 300 ° C. for 1 hour, the oxygen content increased to 0.24 mass%, and when left at 350 ° C. for 1 hour, it increased to 0.59 mass%.
[0200]
On the other hand, the nanocomposite magnetic powder NCP-0 increases the oxygen content only to 0.10% by mass even when left in the atmosphere at 300 ° C. for 1 hour, and contains oxygen after being left at 350 ° C. for 1 hour. The rate is up to 0.20% by mass, and it can be seen that the oxidation resistance is excellent as compared with the conventional quenched magnet powder.
[0201]
Moreover, the result of having measured the heating mass increase rate of each magnetic powder using the thermobalance is shown in FIG. The heating atmosphere was in the air, and the heating rate was 10 ° C./min. As can be seen from FIG. 15, the nanocomposite magnetic powder NCP-0 has less mass increase due to oxidation than the MQP-B and MQP-O, and is excellent in oxidation resistance.
[0202]
Next, with regard to the magnetic characteristics in Table 1, MQP-B has a significant decrease in magnetic characteristics. For example, (BH) max is about 65% of that left at 23 ° C. for 1 hour when left at 300 ° C. for 1 hour. It decreased to about 30% when left at 350 ° C. for 1 hour. Further, the (BH) max of MQP-O decreased to less than about 80% of that left at 350 ° C. for 1 hour after being left at 23 ° C. for 1 hour. On the other hand, even when the nanocomposite magnetic powder NCP-0 was allowed to stand at 350 ° C. for 1 hour, its (BH) max decreased only to about 90% of that left at 23 ° C. for 1 hour.
[0203]
Thus, since the nanocomposite magnetic powder is excellent in oxidation resistance, the magnetic powder is not easily oxidized in the process of producing a bonded magnet (for example, preparation of a compound and / or thermosetting). Therefore, it is possible to simplify or omit the rust preventive treatment of magnetic powder, which is necessary for conventional quenched magnet powder (especially MQP-B) in order to suppress the oxidation of magnetic powder. In addition, a molded body produced by molding a compound needs to be cured by heating, for example, a resin in order to improve its strength. In the case of using the conventional rapidly cooled magnet powder, it was necessary to heat and cure in an inert gas atmosphere such as vacuum or Ar in order to suppress the oxidation of the magnetic powder, but by using the nanocomposite magnetic powder, It becomes possible to carry out heat curing. That is, by using the nanocomposite magnetic powder, the manufacturing process of the bonded magnet can be simplified and the cost can be reduced. Furthermore, since conventional quenching magnet powder has low oxidation resistance, it is applied to a bonded magnet for injection molding that requires a step of kneading with a resin at a temperature of about 250 ° C. to 300 ° C. or a step of molding. However, by using nanocomposite magnetic powder, a bonded magnet produced by injection molding can be obtained. In order to sufficiently obtain the excellent oxidation resistance advantage of the nanocomposite magnetic powder, the oxygen content after being left in the atmosphere at a temperature of 300 ° C. or higher and 350 ° C. or lower for 1 hour is 0.24% by mass or lower. It is preferable to use magnetic powder prepared in the above, and it is preferable to use magnetic powder prepared so that the oxygen content is 0.2% by mass or less. For example, considering the magnetic properties required for bonded magnets for various rotating machines and actuators, the magnetic properties of magnetic powders preferably used for these bonded magnets are as follows: _r ≧ 0.7T, (BH) max ≧ 80 kJ / m ^Three , H _cJ It is preferable that ≧ 600 kA / m is satisfied. When the magnetic powder having the oxidation resistance described above is used, the above magnetic characteristics can be obtained even in consideration of the influence of oxidation in the manufacturing process of the bonded magnet.
[0204]
[Table 1]

[0205]
The nanocomposite magnetic powder according to the present invention has a feature that its magnetic properties are small in particle size dependence due to its composition and structure characteristics. The nanocomposite magnetic powder has a relatively low content of the rare earth element R, and there is no R-rich grain boundary phase. ₂ Fe ₁₄ Small boride phases are dispersed so as to surround the B phase, and since Ti has a high affinity with boron, the boride phase contains more Ti than the other phases. As a result, the nanocomposite magnetic powder is superior in oxidation resistance compared to conventional quenched magnet powder, and can maintain excellent magnetic properties after pulverization.
[0206]
Conventional quenched magnet powders contain a relatively large amount of rare earth element R and are therefore easily oxidized. The smaller the particle size, the more markedly the magnetic properties are degraded due to oxidation of the powder particle surface. For example, in MQP-B (maximum particle size of 300 μm or less), as shown in Table 2, the magnetic properties of powder particles having a particle size of 75 μm or less, particularly 53 μm or less are degraded. Residual magnetic flux density B _r , The residual magnetic flux density B of the powder particles of more than 125 μm and 150 μm or less showing the highest value _r (0.90T), residual magnetic flux density B of powder particles of 53 μm or less _r (0.79T) has fallen to less than 90%. In addition, regarding (BH) max, the average (BH) max (simple average of values of 38 μm or less and more than 38 μm and 53 μm or less) of powder particles of 53 μm or less is 85.5 kJ / m. ^Three 114.6 kJ / m, which is the average (BH) max (simple average of values greater than 150 μm and less than 180 μm and greater than 180 μm and less than 212 μm) for powder particles greater than 150 μm and less than 212 μm ^Three It has fallen to less than 75%.
[0207]
On the other hand, the nanocomposite magnetic powder has a low rate of decrease in magnetic properties due to oxidation, and the particle size dependence of magnetic properties is small. For example, in the nanocomposite magnetic powder NCP-0 (maximum particle size of 300 μm or less), as shown in Table 3, the magnetic properties hardly depend on the particle size and have excellent magnetic properties. For example, residual magnetic flux density B _r Is the residual magnetic flux density B of the powder particles of more than 106 μm and less than 125 μm showing the highest value _r (0.845T), residual magnetic flux density B of powder particles of 53 μm or less _r (About 0.829T) has a value of 98% or more. Also, regarding (BH) max, the average (BH) max of powder particles of 53 μm or less is 104.6 kJ / m. ^Three The average (BH) max of the powder particles of more than 150 μm and less than 212 μm is 106.6 kJ / m ^Three Of 98% or more. As a result of the same evaluation for nanocomposite magnetic powders of various compositions, the average (BH) max of the powder particles of 53 μm or less of the nanocomposite magnetic powders of most compositions is the average (BH) of the powder particles of more than 150 μm and 212 μm or less. ) It has been found that a value of 90% or more of max is obtained, and a value of 95% or more is obtained for many compositions. The particle size distribution of the magnetic powder was evaluated using a standard sieve conforming to JIS8801.
[0208]
[Table 2]

[0209]
[Table 3]

[0210]
Thus, since the nanocomposite magnetic powder has a magnetic property equivalent to or higher than that of the conventional quenched magnet powder, it can be used as a bonded magnet magnetic powder instead of the conventional quenched magnet powder (for example, MQ powder). Of course, the magnetic powder for bonded magnets may be composed only of nanocomposite magnetic powder, but for example, powder particles having a particle size of 53 μm or less in the MQ powder described above may be replaced with nanocomposite magnetic powder.
[0211]
The effect of improving the filling property by mixing fine particles of 53 μm or less and 38 μm or less will be described below by illustrating experimental results.
[0212]
First, samples NCP-1 to NCP-5 of nanocomposite magnetic powders having various particle size distributions as shown in Table 4 were prepared. The NCP-1 magnetic powder was prepared by pulverizing with a power mill using a 0.5 mmφ screen, and the other NCP-2 to NCP-5 magnetic powders were each rotated at the number of revolutions using the pin mill device described above. Was adjusted to 3000 rpm, 4000 rpm, 5000 rpm and 8000 rpm. Table 5 shows the results of measuring the tap density of these magnetic powder samples NCP-1 to NCP-5 using a tap densifier. Table 5 shows the mass% of the powder particles having a particle diameter of 53 μm or less and the mass% of the powder particles having a particle diameter of more than 250 μm included in each magnetic powder sample.
[0213]
As can be seen from the results in Table 5, samples NCP-3 to NCP-5 containing particles having a particle size of 53 μm or less in an amount of 10% by mass or more (strictly 9.5% by mass or more) have a tap density of 4.3 g / cm ^Three It is high as above, and it can be seen that the filling property of the magnetic powder is excellent. The filling property of the magnetic powder evaluated by the tap density of the magnetic powder correlates with the filling property of the compound powder for the bond magnet, and the filling property of the compound powder prepared using the magnetic powder having high filling property is also high. . Therefore, by using the magnetic powder containing 10% by mass of the nanocomposite magnetic powder having a particle size of 53 μm or less, the filling property and fluidity of the bonded magnet compound powder are improved, and a high-quality molded product can be obtained.
[0214]
[Table 4]

[0215]
[Table 5]

[0216]
Furthermore, in order to improve the molding density, it is preferable to include powder particles having a particle size of 38 μm or less. Nanocomposite magnetic powders NCP-11 to NCP-16 having a particle size distribution shown in Table 6 were prepared and mixed with 2% by mass of an epoxy resin to obtain a compound. Molding pressure of 980 MPa (10 t / cm) using each compound ² ) To obtain a bonded magnet molded body. The density of each bonded magnet compact is shown in FIG. 16 together with the content of powder particles having a particle size of 38 μm or less in the magnetic powder used for each compound.
[0217]
[Table 6]

[0218]
As can be seen from FIG. 16, the density of the molded product is lowered when the content of the powder particles of 38 μm or less is too low or too high. As a result of various studies, it is preferable to use magnetic powder containing about 8% by mass or more of powder particles having a particle size of 38 μm or less in order to obtain a sufficient compact density. However, if magnetic powder having a particle size of more than about 16% by mass with a particle size of 38 μm or less is used, the moldability is lowered, and a high-density and high-quality molded product may not be obtained.
[0219]
[Description of Compound and Magnet Manufacturing Method]
The magnetic powder for bonded magnet containing the above-mentioned nanocomposite magnetic powder is mixed with a binder such as resin to produce a compound for bonded magnet.
[0220]
A compound for injection molding is manufactured by kneading magnetic powder and a thermoplastic resin using a known kneading apparatus (for example, a kneader or an extruder). A compound for compression molding is produced by mixing a thermosetting resin diluted with a solvent and magnetic powder and removing the solvent. The obtained mixture of magnetic powder and resin is crushed so as to have a predetermined particle size, if necessary. It is good also as a granular form by adjusting the conditions of crushing. Moreover, you may granulate the powder material obtained by the grinding | pulverization.
[0221]
In order to improve the corrosion resistance of the magnetic powder, a known surface treatment such as chemical conversion treatment may be applied to the surface of the magnetic powder in advance. Furthermore, in order to further improve the corrosion resistance of magnetic powder, wettability with resin, and moldability of compounds, various coupling agents such as silane, titanate, aluminate and zirconate, ceramic ultrafine particles such as colloidal silica, stearin Lubricants such as zinc acid and calcium stearate may be used, and heat stabilizers, flame retardants, plasticizers and the like may be used.
[0222]
Since the compound for magnets is used for various purposes by various molding methods, the type of resin and the blending ratio of the magnetic powder are appropriately determined according to the purpose. Examples of the resin include thermosetting resins such as epoxy resin, phenol resin, and melamine resin, polyamide (nylon 66, nylon 6, nylon 12, etc.), polyethylene, polypropylene, polyvinyl chloride, polyester, polyphenylene sulfide, and the like. Plastic resins, rubber and elastomers, and further modified products, copolymers, and mixtures thereof can be used.
[0223]
Furthermore, since the filling property and / or moldability of the compound is improved by the magnetic powder of the present invention, it is also possible to use a high-viscosity resin that has been difficult to use conventionally. Furthermore, since magnetic powder is difficult to oxidize, it is possible to use a resin that has a high melting point or softening point and could not be used conventionally (for example, polyimide, liquid crystal polymer, and high molecular weight grades of various resins). Magnet properties (heat resistance, etc.) can be improved. Even when a thermosetting resin is used, a resin that cures at a higher temperature than before can be used.
[0224]
Examples of the molding method include compression molding, rolling molding, calendar molding, extrusion molding, and injection molding. Among these molding methods, compression molding, rolling molding, and extrusion molding can form only a relatively simple shaped body, but since a very high fluidity is not required at the time of molding, the filling rate of the magnet powder can be increased. By using the magnetic powder according to the present invention, it is possible to achieve a higher magnetic powder filling rate (for example, more than 80%) than before and to fill up to about 90% at the maximum. However, if the filling rate is increased too much, the resin for sufficiently bonding the magnetic particles is insufficient, and there is a risk that the mechanical strength of the bonded magnet will decrease or the magnetic particles may fall off during use. 85% or less is preferable. Further, by using the magnetic powder of the present invention in compression molding, there is an advantage that the amount of voids formed on the surface of the molded body can be reduced and adverse effects on the resin film formed on the surface can be suppressed. . In these molding methods, a thermosetting resin, a thermoplastic resin, rubber, or the like is appropriately used.
[0225]
Since the fluidity is improved when the magnetic powder according to the present invention is used, it is particularly suitably used for an injection molding compound. It is possible to obtain a molded body having a complicated shape, which has been difficult to be molded by a compound using conventional quenched magnet powder. Moreover, since magnetic powder can be mix | blended with the filling rate (for example, exceeding 65%) higher than before, the magnetic characteristic of a magnet body can be improved. Furthermore, since the magnetic powder according to the present invention has a relatively small content of rare earth elements, it is difficult to be oxidized. Therefore, even if injection molding is performed at a relatively high temperature using a thermoplastic resin or thermoplastic elastomer having a relatively high softening point, the magnetic properties are not deteriorated.
[0226]
[Application example of bonded magnet]
As described above, the bonded magnet compound according to the present invention has an excellent filling property (formability) as compared with a compound using a conventional quenched magnet powder (for example, product name MQP-B manufactured by MQI). Since it is possible to form a bonded magnet having excellent heat resistance and having a magnetic property equivalent to or higher than that of a conventional bonded magnet using a quenched magnet powder, it is suitably used for various applications.
[0227]
An example applied to a stepping motor will be described with reference to FIG.
[0228]
FIG. 17 is an exploded perspective view schematically showing a configuration of a stepping motor 100 having a permanent magnet rotor type. The stepping motor 100 includes a rotor 101 and a stator portion 102 provided around the rotor 101. The rotor 101 includes a bonded magnet in which an outer peripheral surface having an outer diameter of 8 mm is evenly magnetized to 10 poles. The stator unit 102 includes outer yokes 102a and 102b, two inner yokes 103 joined to each other back to back, and exciting coils 104a and 104b accommodated therebetween. The stepping motor 100 is a so-called PM type pulse motor that performs an operation in which the rotor 101 is displaced by one step angle by the magnetomotive force of the exciting coils 104a and 104b corresponding to one pulse current.
[0229]
The bonded magnet included in the rotor 101 is formed using the above-described compound having excellent filling properties (formability) according to the present invention, and has a magnetic property equivalent to or higher than that of a conventional bonded magnet using a rapidly cooled magnet powder. Excellent mechanical properties, no risk of chipping, and excellent reliability. Moreover, it is excellent also in heat resistance.
[0230]
A stepping motor having a bonded magnet formed using the compound according to the present invention has a small size, high performance and excellent reliability, and is an OA device such as a printer or a disk drive device, or an AV such as a camera or a video. It is suitably used for equipment.
[0231]
The rotor 101 can be manufactured by various methods. For example, it may be formed by compression molding a compound using a thermosetting resin, or may be formed by injection molding or extrusion molding a compound using a thermoplastic resin. Hereinafter, a method for manufacturing the rotor 101 will be described with reference to FIG.
[0232]
For example, in the case of using a compound having a thermosetting resin as a binder, a bonded magnet shown in FIG. 18 (d) is adopted by adopting a molding method as described with reference to FIGS. 18 (a) to (c). The integrally molded rotor 200 can be manufactured.
[0233]
A rotor 200 shown in FIG. 18D includes a rotor shaft 205, a yoke 208 provided around the rotor shaft 205, and a bond magnet 210. Bond magnet 210 is bonded to surface 212 of yoke 208.
[0234]
The rotor 200 is manufactured by the steps shown in FIGS.
[0235]
First, as shown in FIG. 18A, first, the compound box 201 is filled into the cavity 202 while sliding the feeder box 203 containing the powdery compound 201 on the upper surface of the mold 204. The mold 204 is set with a yoke 208 into which the rotor shaft 205 is press-fitted in the center, and an auxiliary member 207 is provided so as to cover the rotor shaft 205. A cavity 202 is formed between the mold 204 and them.
[0236]
Next, as shown in FIG. 18B, for example, the compound 201 is compression-molded through the upper punch 209, so that the yoke 208 and the molded body of the compound 201 are physically coupled.
[0237]
Next, as shown in FIG. 18C, the rotor molded body is taken out from the mold 204. The auxiliary member 207 is easily removed from the rotor shaft 205 and the yoke 208, and the rotor shaft 205, yoke 208, and bond magnet 210 are integrated. However, in this state, the bond magnet 210 is a compound powder compact, and the thermosetting resin contained in the compound is uncured.
[0238]
Next, the compound is cured at a predetermined temperature in order to cure the bond magnet 210 and to strengthen the bond at the interface 212 between the yoke 208 and the bond magnet 210. The curing temperature and the curing time are appropriately set according to the resin used.
[0239]
Since the compound by this invention contains the magnetic powder which is excellent in heat resistance, it can be a compound hardened | cured suitably at the curing temperature higher than before. Therefore, it is possible to form the bonded magnet 210 that is superior in heat resistance, mechanical properties, and adhesive strength as compared with the conventional case. Furthermore, since the compound according to the present invention is excellent in corrosion resistance, the magnetic powder itself has very little deterioration in magnet characteristics even when the thermosetting treatment is performed in the air. Therefore, it is not necessary to perform the thermosetting treatment in an inert atmosphere, so that process costs can be reduced.
[0240]
According to the molding method described above, the yoke 208, the rotor shaft 205, and the bond magnet 210 can be molded integrally while molding the ring-shaped bond magnet 210, so that the rotor 200 can be manufactured with high productivity. Can do.
[0241]
In addition, although the example which hardens | cures after taking out a molded object from the metal mold | die 204 was demonstrated, even if it hardens | cures in the metal mold | die 204 by providing a heating mechanism in the metal mold | die 204, even if it hardens | cures in the pressurized state Good. Furthermore, the present invention is not limited to compression molding, and a bonded magnet integral molding rotor can be formed by injection molding.
[0242]
Further, the compound according to the present invention has a high filling property (formability and / or fluidity) as compared with the compound using the conventional quenched magnet powder, so that it is surely filled in a small gap (for example, about 2 mm width). be able to. Therefore, the compound according to the present invention is suitably used for manufacturing a magnet-embedded rotor 300 (see FIG. 19) used in an IPM (Interior Permanent Magnet) type motor.
[0243]
A magnet-embedded rotor 300 shown in FIG. 19 includes an iron core (for example, a diameter of 80 mm, a thickness of 50 mm) 301, a rotation shaft slot 302 formed at the center of the iron core 301, and a plurality of parts formed around the iron core 301. Arc-shaped magnet slot 304. Here, eight arc-shaped magnet slots 304 are provided, and each slot 304 has a first slot (for example, a width of 3.5 mm) 304a and a second slot (for example, a width of 1.8 mm) 304b. It has a layer structure. These slots 304a and 304b are filled with a compound according to the present invention to form a bonded magnet. An IPM motor is obtained by combining with a stator (not shown) in which S poles and N poles are alternately arranged so as to face a plurality of magnet slots 304 of the rotor 300.
[0244]
The formation of the bonded magnet can be performed in various ways. For example, when a compound containing a thermosetting resin is used, an in-slot compression molding method (see, for example, JP-A-63-98108) can be employed. Further, when a compound containing a thermoplastic resin is used, an extrusion molding method or an injection molding method can be employed. Whichever molding method is employed, the compound according to the present invention is excellent in filling properties, so that it can be reliably filled into the slots 304a and 304b, has excellent mechanical properties and heat resistance, and has a magnetic property equivalent to or higher than that of the conventional method. A bonded magnet having characteristics can be formed. Therefore, it is possible to manufacture a small IPM type motor having higher performance and higher reliability than conventional ones.
[0245]
The compound by this invention is used suitably for formation of the bond magnet which the angle sensor (rotary encoder) 400 shown to Fig.20 (a) has.
[0246]
A rotary encoder 411 shown in FIG. 20A includes a rotating shaft 413, a rotating drum 416 connected to the rotating shaft 413, and a plurality of bond magnets 415 bonded to the outer peripheral surface of the rotating drum 416. And a detector 417 arranged on the outer peripheral surface of the rotor 414 so as to be separated from each other. The detector 417 is not particularly limited as long as it can detect a change in magnetic flux from the rotor 414. For example, a Hall element, a magnetoresistive element, or a magnetoimpedance effect element can be used. The rotating shaft 413 is connected to the motor 412. The detector 17 is connected to a measurement unit (not shown).
[0247]
The bonded magnet 415 formed using the compound according to the present invention has a cylindrical shape as shown in FIG. 20B, for example, and has an N pole and an S pole along the outer peripheral surface of the rotating drum 416. Alternatingly arranged. The bond magnet 415 and the rotating drum 416 are joined by, for example, an adhesive. The rotating drum 416 is formed using, for example, a metal material and may not be a magnetic material.
[0248]
The rotary encoder 400 operates as follows. When the rotating shaft 413 of the motor 412 rotates, the rotor 414 rotates according to the rotation. At this time, the direction of the magnetic flux formed in the detector 417 by the bonded magnet 415 disposed on the outer peripheral surface of the rotor 414 changes according to the rotation of the rotor 414. The detector 417 generates an output signal (such as a voltage change amount or a current change amount) corresponding to the change in the direction of the magnetic flux and outputs the output signal to a measurement unit (not shown). In this way, the rotation amount (angle) of the motor 412 is measured.
[0249]
The compound according to the present invention is excellent in filling properties (formability, fluidity), has a magnetic property equal to or higher than conventional, and can form a bonded magnet that is superior in mechanical properties and heat resistance than conventional, A small, high-performance and highly reliable angle sensor can be manufactured.
[0250]
Furthermore, the compound by this invention is used suitably for formation of the bond magnet for magnetic rolls demonstrated referring FIG. 21 (a) and (b).
[0251]
FIG. 21A is a cross-sectional view schematically showing the structure of an electrophotographic process cartridge 501. The cartridge 501 integrally includes a photosensitive drum 510 that is rotationally driven in the direction of arrow A, a charging roller 502 that charges the photosensitive drum 510, a developing device 511, and a cleaning device 512.
[0252]
The developing device 511 includes a developing container 509 that stores toner 513, and a developing sleeve 506 is rotatably disposed at an opening of the developing container 509 so as to face the photosensitive drum 510. Further, the developing device 511 includes an elastic blade 514, and the elastic blade 514 contacts the developing sleeve 506 and regulates the layer thickness of the toner 513 carried and conveyed by the developing sleeve 506.
[0253]
FIG. 21B is a cross-sectional view schematically showing the configuration of the developing device 511 included in the process cartridge 501.
[0254]
The developing sleeve 506 is made of a nonmagnetic material and is rotatably supported by the developing container 509 via a bearing. A magnetic roll (for example, 8.5 mm in diameter) 507 is disposed in the developing sleeve (for example, 10 mm in diameter) 506. A notch 507a-1 is formed in the shaft portion 507a of the magnetic roll 507, and the magnetic roll 507 is fixed by engaging the notch 507a-1 with the developing container 509. The magnetic roll 507 has a developing pole S1 at a position facing the photosensitive drum 510, and has an S2, N1, and N2 poles at other positions.
[0255]
The magnet 508 is disposed so as to surround the developing sleeve 506. A magnetic curtain is formed in the gap g with the developing sleeve 506, and toner is held in the gap by the magnetic curtain, thereby preventing toner leakage. The
[0256]
Since the magnetic roll 507 is formed using the compound according to the present invention, the magnetic roll 507 has a magnetic property equivalent to or higher than that of a conventional magnet, and is excellent in mechanical properties and heat resistance. Therefore, it is possible to further reduce the size of the magnetic roll 507 and the developing sleeve 506 as compared with the conventional one, and to improve the performance. The magnetic roll formed by using the compound according to the present invention can be applied to a developing device and a developing cartridge in a copying machine or a laser beam printer.
[0257]
【Example】
(Example 1)
B, Fe, Ti, Nd, and Nb having a composition shown in Table 7 below and having a purity of 99.5% or more were weighed so as to have a total amount of 600 grams (g), and charged into an alumina crucible. Thereafter, these alloy raw materials were melted in an argon (Ar) atmosphere at a pressure of 70 kPa by high-frequency heating to produce a molten alloy. After the molten metal temperature reached 1500 ° C., it was cast on a water-cooled copper mold to produce a flat alloy.
[0258]
After the obtained alloy was pulverized, 25 milligrams (mg) of crushed pieces were dissolved, and the solidification process of the molten alloy was analyzed at a cooling rate of 20 ° C./min using a differential calorimeter (DTA) in an Ar stream. Table 7 shows the measurement results.
[0259]
[Table 7]

[0260]
Here, Sample No. Nos. 1 to 5 are samples in which Nb is added in addition to Ti. 6 to 7 are samples to which Nb was not added.
[0261]
In the rightmost column of Table 7, each sample No. About the alloy melts 1-7, the temperature characterizing the solidification process of the alloy melt is described. The temperature described as “1st” indicates the temperature at which the first solidification occurs in the process of cooling the molten alloy (“liquidus temperature”). The temperature described as “2nd” indicates the temperature at which the next solidification occurs at a temperature lower than the liquidus temperature in the process of cooling the molten alloy (“freezing point”). Specifically, these temperatures are temperatures at which an exothermic peak is observed by a differential calorimeter (DTA).
[0262]
FIG. 2 (Nb addition) and Sample No. It is a graph which shows DTA of 6 (Nb additive-free). As is clear from FIG. In the case of 2, sample No. 6, the temperature of the first exothermic peak generated in the cooling process, that is, the liquidus temperature (“1st”) is lowered by 60 ° C. or more.
[0263]
This first exothermic peak is TiB ₂ This may be caused by the precipitation of a compound phase of titanium and boron. In this example, since Ti and boron are added at a higher concentration than before, a compound of titanium and boron (high melting point) is easily formed, and the deposition temperature is estimated to be high. Conventional composition system without adding Ti (Fe _Three B / Nd ₂ Fe ₁₄ In the B system), the liquidus temperature of the molten alloy was about 1200 ° C. or less. In the present invention, it is considered that by adding Nb together with Ti, the precipitation temperature of such a compound is lowered and the liquidus temperature of the molten alloy is lowered.
[0264]
Sample No. 6 (Comparative Example), it is necessary to perform strip casting at a high tapping temperature of about 1350 ° C. When the alloy of No. 2 (Example) is used, it is possible to set the tapping temperature to about 1250 ° C., for example. When the temperature of the hot water is reduced in this way, R that precipitates early in the cooling process of the molten metal. ₂ Fe ₁₄ B-type compounds and TiB ₂ Is suppressed, and the magnet characteristics are improved.
[0265]
(Example 2)
B, Fe, Ti, Nd, and C having a composition shown in Table 2 below and having a purity of 99.5% or more were weighed so as to have a total amount of 600 g, and charged into an alumina crucible. Thereafter, these alloy raw materials were melted in an argon (Ar) atmosphere at a pressure of 70 kPa by high-frequency heating to produce a molten alloy. After the molten metal temperature reached 1500 ° C., it was cast on a water-cooled copper mold to produce a flat alloy.
[0266]
After the obtained alloy was pulverized, 25 milligrams (mg) of crushed pieces were dissolved, and the solidification process of the molten alloy was analyzed at a cooling rate of 20 ° C./min using a differential calorimeter (DTA) in an Ar stream. . Table 8 shows the measurement results.
[0267]
[Table 8]

[0268]
Here, Sample No. 8 to 13 are samples in which C is added together with Ti. 14 to 15 are samples to which C was not added.
[0269]
In the rightmost column of Table 8, each sample No. About the alloy melts of 8-15, the temperature which characterizes the solidification process of an alloy melt is described. The temperature described as “1st” indicates the temperature at which the first solidification occurs in the process of cooling the molten alloy (“liquidus temperature”). The temperature described as “2nd” indicates the temperature at which the next solidification occurs at a temperature lower than the liquidus temperature in the process of cooling the molten alloy (“freezing point”). Specifically, these temperatures are temperatures at which an exothermic peak is observed by a differential calorimeter (DTA).
[0270]
FIG. 8 (C addition) and Sample No. It is a graph which shows DTA of 14 (C no addition). As is clear from FIG. In the case of 8, sample no. 14, the temperature of the first exothermic peak generated in the cooling process, that is, the liquidus temperature (“1st”) is lowered by about 40 ° C.
[0271]
This first exothermic peak is TiB ₂ This may be caused by the precipitation of a compound phase of titanium and boron. In this example, since Ti and boron are added at a higher concentration than before, a compound of titanium and boron (high melting point) is easily formed, and the deposition temperature is estimated to be high. Conventional composition system without adding Ti (Fe _Three B / Nd ₂ Fe ₁₄ In the B system), the liquidus temperature of the molten alloy was about 1200 ° C. or less. In the examples of the present invention, it was considered that the addition of C together with Ti lowered the precipitation temperature of such a compound and lowered the liquidus temperature of the molten alloy.
[0272]
Sample No. When the alloy No. 14 is used, strip casting or the like needs to be performed at a high tapping temperature of about 1350 ° C. When using the alloy of No. 8 (Example), it is possible to set the tapping temperature to about 1300 ° C., for example. When the temperature of the hot water is reduced in this way, R that precipitates early in the cooling process of the molten metal. ₂ Fe ₁₄ B-type compounds and TiB ₂ Is suppressed, and the magnet characteristics are improved.
[0273]
Next, B, Fe, Ti, Nd, and C having a composition shown in Table 8 and having a purity of 99.5% or more are weighed so that the total amount is 15 g, and an orifice having a diameter of 0.8 mm is provided at the bottom. It put into the quartz crucible. Thereafter, these alloy raw materials were melted in an Ar atmosphere at a pressure of 1.33 to 47.92 kPa by high-frequency heating to produce a molten alloy. After the molten metal temperature reached 1350 ° C., the molten metal surface was pressurized with Ar gas, and the molten metal was dropped from the orifice onto the outer peripheral surface of the cooling roll located 0.7 mm below. The cooling roll was made of pure copper and was rotated so that the outer peripheral surface speed was 15 m / sec. Due to such contact with the cooling roll, the molten alloy was rapidly cooled and solidified. Thus, a continuous rapidly solidified alloy ribbon having a width of 2 to 3 mm and a thickness of 20 to 50 μm was obtained. FIG. 8 and sample no. It is a graph which shows the pattern of 14 XRD. As is clear from FIG. In the case of No. 8, amorphous material occupies the majority, whereas sample No. In No. 14, the ratio of the crystal structure is large.
[0274]
The rapidly solidified alloy ribbon was held in a heat treatment temperature range of 600 to 800 ° C. in an Ar atmosphere for 6 to 8 minutes, and then cooled to room temperature. Thereafter, the magnetic properties of the quenched alloy ribbon (length 3 to 5 mm) were evaluated using VSM. Table 9 shows the measurement results.
[0275]
[Table 9]

[0276]
Next, Sample No. A raw material alloy having the same composition as that of No. 13 was prepared, and a quenched alloy was prepared using a strip casting apparatus as shown in FIG. Specifically, B, Fe, Ti, Nd, and C having a purity of 99.5% or more were weighed so that the total amount was 10 kg, and charged into the dissolution tank. Thereafter, these alloy raw materials were melted in an Ar atmosphere at a pressure of 30 kPa by high-frequency heating to prepare a molten alloy. After the molten metal temperature reached 1350 ° C., the molten metal was poured into the chute. The molten metal flowed smoothly on the chute and was cooled by a cooling roll. The surface peripheral speed of the cooling roll was 12 m / sec.
[0277]
The quenched alloy thus obtained (average thickness: about 80 μm) was held in an Ar atmosphere at a heat treatment temperature range of 740 ° C. for 6 to 8 minutes, and then cooled to room temperature. Thereafter, the magnetic properties of the quenched alloy were evaluated using VSM.
[0278]
The measurement result is the residual magnetic flux density B _r Is 0.79T, coercive force H _cJ 1090 kA / m, maximum magnetic energy product (BH) _max Is 102kJ / m ^Three Met. The magnetic characteristics are shown in Sample No. As compared with the magnetic characteristics of FIG. 8, it can be seen that almost the same characteristics were obtained.
[0279]
Next, XRD and a demagnetization curve were measured for the sample in which the ratio of C to the total of (B + C) (atomic ratio p) was 0.25 or less and the sample in which p exceeded 0.25.
[0280]
FIG. 11 shows Nd ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Four (Example: p = 0.1) and Nd ₉ Fe ₇₃ B ₇ C ₇ Ti _Four The XRD pattern before the heat processing of (Comparative example: p = 0.5) is shown. Although these samples differed in composition, they were produced in the same manner as in the above-described examples. In FIG. 12, Nd ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Four (Example) and Nd ₉ Fe ₇₃ B ₇ C ₇ Ti _Four The demagnetization curve of (comparative example) is shown.
[0281]
When the C ratio p exceeds 0.25 and reaches 0.5, as shown in FIG. 11, a Ti—C diffraction peak is remarkably observed. When there is too much C in this manner, a large amount of Ti-C phase precipitates in the quenched alloy, so the constituent phase ratio after heat treatment deviates from the desired range, and the squareness of the demagnetization curve becomes worse as shown in FIG. . If the ratio of C to the total of (B + C) (atomic ratio p) was 0.25 or less, such a problem did not occur.
[0282]
(Example 3)
In this embodiment, the strip casting apparatus shown in FIG. 3 was used.
[0283]
First, Nd by atomic ratio ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Three Nb ₁ So that the total amount is 5 kg using metals of B, C, Fe, Nb, Ti, and Nd having a purity of 99.5% or more. These metals were put into an alumina crucible and dissolved by high frequency heating in an argon atmosphere at a pressure of 35 kPa. The dissolution temperature was 1350 ° C.
[0284]
After melting, the crucible was tilted and the molten metal was supplied onto a porous ceramic chute and led to the surface of the cooling roll. The surface temperature of the chute was kept at 600 ° C. by a heater. Further, the chute was tilted by 20 ° (= angle α) with respect to the horizontal direction so that the molten metal smoothly flows toward the roll on the chute. Further, the chute was arranged so that the molten metal was poured from the position directly above the roll to the position of the crucible at a position inclined by 40 ° (= angle β). In addition, the chute | shoot in a present Example has the molten metal guide for dividing the flow of the molten metal received from the crucible into two strips, and supplying to a roll, as shown in FIG.
[0285]
The cooling roll was rotated at a surface peripheral speed of 14 m / sec. By adjusting the tilt angle of the crucible, the supply rate of the molten metal flowing on the chute was adjusted to 1.5 kg / min per one flow of the molten metal. In this embodiment, a pure copper roll having a surface center line roughness Ra of 5 μm was used. The roll temperature was prevented from rising by water cooling inside the roll.
[0286]
When the structure of the obtained quenched alloy was examined by the characteristic X-ray of CuKα, Nd ₂ Fe ₁₄ Along with the diffraction peak of B, Fe _{twenty three} B ₆ It was confirmed that the quenched alloy structure was mixed with α-Fe.
[0287]
FIG. 13 shows the obtained rapidly cooled alloy powder XRD, and FIG. 14 shows a demagnetization curve of the quenched alloy measured using a vibration magnetometer. In FIG. 13 and FIG. 14, the curve described as “as-cast” relates to the quenched alloy.
[0288]
Next, the quenched alloy was pulverized by a power mill. Thereafter, argon gas was flowed, and the quenched alloy powder was supplied into a hoop belt type continuous heat treatment furnace maintained at 740 ° C. to perform heat treatment. At this time, the powder feeding rate was kept at 30 g / min.
[0289]
The powder XRD and demagnetization curve after heat treatment are also shown in FIGS. 13 and 14, respectively. In FIG. 13 and FIG. 14, the data after the heat treatment is shown by a curve described as “as-annealed”. The magnetic properties after the heat treatment are shown in Table 10 below.
[0290]
[Table 10]

[0291]
As can be seen from FIG. 14 and Table 10, the iron-based permanent magnet in this example exhibited good magnetic properties.
[0292]
Next, the fine metal structure after the heat treatment was observed with a transmission electron microscope (TEM). As a result, it was found that crystal grains having an average grain size of about 40 nm and fine crystal grains of about 10 nm exist at the grain boundaries in the structure after the heat treatment. As a result of metallographic analysis by HRTEM (High Resolution Transmission Electron Microscope), crystal grains having an average grain size of about 40 nm are Nd. ₂ Fe ₁₄ B, and the grain boundary is Fe _{twenty three} B ₆ Or Fe _Three It was confirmed that iron-based boride of B was present.
[0293]
Example 4
Also in this example, the strip casting apparatus shown in FIG. 3 was used.
[0294]
First, Nd by atomic ratio ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Three Nb ₁ So that the total amount is 5 kg using metals of B, C, Fe, Nb, Ti, and Nd having a purity of 99.5% or more. These metals were put into an alumina crucible and dissolved by high-frequency heating in an argon atmosphere at a pressure of 35 kPa. The dissolution temperature was 1350 ° C.
[0295]
After melting, the crucible was tilted, and the molten metal was supplied onto a porous ceramic chute and led to the surface of the cooling roll. The surface temperature of the chute was kept at 600 ° C. by a heater. Further, the chute was tilted by 20 ° (= angle α) with respect to the horizontal direction so that the molten metal smoothly flows toward the roll on the chute. Moreover, the chute | shoot was arrange | positioned so that molten metal might be poured into the position inclined only 40 degrees (= angle (beta)) from the direct upper part of a roll to the position of a crucible. Also in this example, the chute shown in FIG. 4 was used.
[0296]
In this example, the cooling roll was rotated at the surface peripheral speed shown in Table 11. Further, by adjusting the tilt angle of the crucible, the supply speed (per one flow) of the molten metal flowing on the chute was adjusted as shown in Table 11. The width of one flow of the molten metal was 10 mm, and the influence of the roll peripheral speed and the molten metal supply speed on the rapid cooling was examined.
[0297]
In this example, as in Example 3, a pure copper roll having a surface centerline roughness Ra of 5 μm was used. The roll temperature was prevented from rising by water cooling inside the roll.
[0298]
[Table 11]

[0299]
In Table 11, “◯” indicates a case where a quenched alloy could be stably produced. On the other hand, “x” indicates a case where splash occurs and a quenched alloy having a desired structure cannot be stably obtained. “Δ” indicates a case in which the production of a stable quenched alloy was frequently observed, but the splash occurred intermittently.
[0300]
From Table 11, when the roll surface peripheral speed is 10 m / second or more and 18 m / second or less, stable rapid cooling is realized when the molten metal supply speed per molten metal flow is 1.0 kg / minute or more and 2.0 kg / minute or less. I understand that. The faster the roll surface peripheral speed, the thinner the quenched alloy ribbon and the more likely the splash will occur.
[0301]
The molten metal supply speed per one flow of the molten metal does not significantly affect the thickness of the quenched alloy ribbon, but changes the width of the quenched alloy ribbon. The larger the melt supply rate, the wider the width of the quenched alloy ribbon.
[0302]
The thickness of the quenched alloy ribbon varies depending on the roll surface peripheral speed. That is, the faster the roll surface peripheral speed, the thinner the quenched alloy ribbon. For example, when the roll surface circumferential speed is 10 m / sec, the average thickness of the quenched alloy ribbon is about 100 μm, and when the roll surface circumferential speed is 22 m / sec, the average thickness of the quenched alloy ribbon is 45-80 μm. Degree.
[0303]
As described above, the thicker the quenched alloy ribbon is (for example, when the thickness is greater than 80 μm), the easier it is to obtain powder particles having a shape close to an equiaxed shape by grinding the quenched alloy. If a bonded magnet is produced using powder containing a large number of particles having an aspect ratio close to 1, a bonded magnet having excellent magnet characteristics can be obtained.
[0304]
In addition, the structure of the quenched alloy prepared under the conditions of the surface peripheral speed of the roll of 14 m / sec and the supply amount of the molten metal per one flow of the molten metal of 1.3 kg / min was examined by the characteristic X-ray of CuKα. As a result, Nd ₂ Fe ₁₄ Along with the diffraction peak of B, Fe _{twenty three} B ₆ It was confirmed that the quenched alloy structure was mixed with α-Fe.
[0305]
【The invention's effect】
According to the present invention, by using a strip cast method and quenching a molten rare earth alloy in which Ti is added to a raw material alloy, the coercive force and magnetization are sufficiently high while reducing the amount of rare earth elements required for the magnet. It is possible to mass produce raw material alloys for iron-based rare earth magnets that exhibit excellent magnetic properties.
[Brief description of the drawings]
FIG. 1 is a diagram showing a configuration example of a strip cast apparatus preferably used in the present invention.
FIG. 2 is a view showing a structure of a nanocomposite magnet manufactured according to the present invention.
FIG. 3 is a diagram showing another configuration example of a strip cast apparatus preferably used in the present invention.
FIG. 4 is a perspective view showing a molten alloy chute (guide means) used in a strip casting apparatus.
FIG. 5 is a diagram showing the influence of the center line roughness Ra on the surface of the cooling roll used in the strip casting method on the rapid cooling of the molten alloy.
FIG. 6 is a diagram showing the influence of the center line roughness Ra on the surface of the cooling roll used in the melt spinning method on the rapid cooling of the molten alloy.
FIG. 7 is a cross-sectional view showing a microstructure of a quenched alloy formed by a strip cast method, where (a) shows a cross section of an RTB-based alloy to which Ti is added, and (b) shows the addition of Ti. The cross section of the conventional R-T-B type alloy which does not perform is shown.
FIG. 2 and Sample No. 6 is a graph showing a DTA of 6.
FIG. 8 and sample no. 14 is a graph showing 14 DTAs.
FIG. 10 shows a sample No. before crystallization heat treatment (as-cast). 8 and sample no. 14 is a graph showing 14 powder X-ray diffraction data.
FIG. 11: Nd ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Four (Example: p = 0.1) and Nd ₉ Fe ₇₃ B ₇ C ₇ Ti _Four The XRD pattern before the heat processing of (Comparative example: p = 0.5) is shown.
FIG. 12 Nd ₉ Fe ₇₃ B _12.6 C _1.4 Ti _Four (Example) and Nd ₉ Fe ₇₃ B ₇ C ₇ Ti _Four The demagnetization curve of (comparative example) is shown.
FIG. 13 is a graph of powder XRD for an example of the present invention. The curve labeled “as-cast” relates to the quenched alloy and the curve labeled “as-annealed” relates to the alloy after heat treatment.
FIG. 14 is a graph of a demagnetization curve relating to an example of the present invention measured using a vibration type magnetometer. The curve described as “as-cast” relates to the quenched alloy, and the curve described as “as-annealed” relates to the alloy after heat treatment.
FIG. 15 is a graph showing the heating mass increase rate of the nanocomposite magnetic powder according to the present invention and the conventional quenched magnet powder.
FIG. 16 is a graph showing the density of a bonded magnet molded body formed using nanocomposite magnetic powders having different particle size distributions.
FIG. 17 is an exploded perspective view schematically showing a configuration of a stepping motor 100 including a permanent magnet rotor type according to an embodiment of the present invention.
FIGS. 18A to 18D are views showing a bonded magnet-integrated mold rotor 200 according to an embodiment of the present invention and a molding process thereof;
FIG. 19 is a schematic diagram showing the structure of a magnet-buried rotor 300 according to an embodiment of the present invention.
FIGS. 20A and 20B are views schematically showing the structure of a rotary encoder 411 according to an embodiment of the present invention.
FIGS. 21A and 21B are cross-sectional views schematically showing the structure of an electrophotographic process cartridge 501 including a magnetic roll 507 according to an embodiment of the present invention.
[Explanation of symbols]
1 Melting furnace
2 Bottom exit of melting furnace
3 molten alloy
4 樋
5 Chute (Melt guide method)
6 Alloy molten metal paddle
7 Cooling roll
8 Quenched alloy

Claims (12)

Composition formula _{(Fe 1-m T m)} 100-xyzn Q x R y Ti z M n (T is at least one element selected from the group consisting of Co and Ni, Q is a group consisting of B and C One or more selected elements, which always contain B, R is a rare earth metal element, M is Al, Si, V, Cr, Mn, Zn, Cu, Ga, Zr, Nb, Mo, Hf, Ta , W, Pt, Pb, Au, and Ag, one or more elements selected from the group consisting of, and composition ratios (atomic ratios) x, y, z, m, and n are 10 <x ≦ A quenched alloy satisfying 20 atomic%, 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦ 5 atomic%, There are crystal structures on two end faces that are 50 μm or more and 200 μm or less and whose width is within a range of 5 mm or more and 20 mm or less and orthogonal to the thickness direction. The crystal structure includes a ferromagnetic iron-based boride phase having an average particle size of 1 nm to 50 nm and an R ₂ Fe ₁₄ B type compound phase having an average particle size of 20 nm to 200 nm, and the ferromagnetic structure iron-based boride phase, the present in grain boundaries or sub-boundaries of the R ₂ Fe ₁₄ B type compound phase, and the rapidly solidified alloy abundance ratio of the R ₂ Fe ₁₄ B type compound phase is 60 vol% or more Nanocomposite magnet powder obtained by crushing,
A compound for a bonded magnet comprising a binder.   The compound for bonded magnets according to claim 1, wherein the binder is a thermosetting resin or a thermoplastic resin.   The said nanocomposite magnet powder is a compound for bonded magnets of Claim 1 or 2 containing 10 mass% or more of powder particles with a particle size of 53 micrometers or less. Composition formula _{(Fe 1-m T m)} 100-xyzn Q x R y Ti z M n (T is at least one element selected from the group consisting of Co and Ni, Q is a group consisting of B and C One or more selected elements, which always contain B, R is a rare earth metal element, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Hf, Ta 1 or more elements selected from the group consisting of W, Pt, Pb, Au and Ag), and the composition ratios x, y, z, m and n are 10 <x ≦ 20 atomic%, Nanocomposite magnet powder satisfying 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦ 5 atomic%, The diameter is in the range of 10 μm to 300 μm, and the ratio of the short axis size to the long axis size is 0.3 or less. The crystal structure in the magnet powder is a ferromagnetic iron-based boride phase having an average particle size of 1 nm to 50 nm and an R ₂ Fe ₁₄ B type compound phase having an average particle size of 20 nm to 200 nm. wherein, the ferromagnetic iron-based boride phase, the present in grain boundaries or sub-boundaries of the R ₂ Fe ₁₄ B type compound phase, and the R ₂ Fe ₁₄ existence ratio of B type compound phase 60 vol% Nanocomposite magnet powder that is
A compound for a bonded magnet comprising a binder.   The bonded magnet compound according to claim 4, wherein the binder is a thermosetting resin or a thermoplastic resin.   The said nanocomposite magnet powder is a compound for bond magnets of Claim 4 or 5 containing 10 mass% or more of powder particles with a particle size of 53 micrometers or less. A method for producing a compound for a bonded magnet comprising a nanocomposite magnet powder and a binder,
Composition formula _{(Fe 1-m T m)} 100-xyzn Q x R y Ti z M n (T is at least one element selected from the group consisting of Co and Ni, Q is a group consisting of B and C One or more selected elements, which necessarily contain B, R is a rare earth metal element, M is Al, Si, V, Cr, Mn, Zn, Cu, Ga, Zr, Nb, Mo, Hf, Ta , W, Pt, Pb, Au, and Ag, one or more elements selected from the group consisting of, and the composition ratio (atomic ratio) x, y, z, m, and n are 10 <x ≦ A quenched alloy satisfying 20 atomic%, 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦ 5 atomic%, and having a thickness Is in the range of 50 to 200 μm, the width is in the range of 5 to 20 mm, and the crystal pair is formed on two end faces perpendicular to the thickness direction. Woven is formed, the crystal structure may include a mean particle size of less ferromagnetic iron-based boride phase 50nm or 1 nm, the average particle diameter of 20nm or 200nm following an R ₂ Fe _{1 4} B compound phase, the ferromagnetic iron-based boride phase exists in the grain boundary or sub-boundaries of the R ₂ Fe ₁₄ B compound phase, and the abundance ratio of the R ₂ Fe ₁₄ B type compound phase is 60 vol% or more Preparing a quenched alloy; and
Producing a nanocomposite magnet powder by pulverizing the quenched alloy; and
The manufacturing method of the compound for bonded magnets including the process of mixing the said nanocomposite magnet powder and binder.   The said binder is a manufacturing method of the compound for bonded magnets of Claim 7 which is a thermosetting resin or a thermoplastic resin.   The said nanocomposite magnet powder is a manufacturing method of the compound for bonded magnets of Claim 7 or 8 containing 10 mass% or more of powder particles with a particle size of 53 micrometers or less. Composition formula _{(Fe 1-m T m)} 100-xyzn Q x R y Ti z M n (T is at least one element selected from the group consisting of Co and Ni, Q is a group consisting of B and C One or more selected elements, which always contain B, R is a rare earth metal element, M is Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Hf, Ta 1 or more elements selected from the group consisting of W, Pt, Pb, Au and Ag), and the composition ratios x, y, z, m and n are 10 <x ≦ 20 atomic%, Nanocomposite magnet powder satisfying 6 ≦ y <10 atomic%, 0.5 ≦ z ≦ 6 atomic%, 0 ≦ m ≦ 0.5, and 0 ≦ n ≦ 5 atomic%, The diameter is in the range of 10 μm to 300 μm, and the ratio of the short axis size to the long axis size is 0.3 or less. The crystal structure in the magnet powder is a ferromagnetic iron-based boride phase having an average particle size of 1 nm to 50 nm and an R ₂ Fe ₁₄ B type compound phase having an average particle size of 20 nm to 200 nm. wherein, the ferromagnetic iron-based boride phase, the present in grain boundaries or sub-boundaries of the R ₂ Fe ₁₄ B type compound phase, and the R ₂ Fe ₁₄ existence ratio of B type compound phase 60 vol% A step of preparing the nanocomposite magnet powder as described above;
The manufacturing method of the compound for bonded magnets including the process of mixing the said nanocomposite magnet powder and binder.   The manufacturing method of the compound for bond magnets of Claim 10 whose said binder is a thermosetting resin or a thermoplastic resin.   The said nanocomposite magnet powder is a manufacturing method of the compound for bonded magnets of Claim 10 or 11 containing 10 mass% or more of powder particles with a particle size of 53 micrometers or less.

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