JP3710837B2

JP3710837B2 - Rare earth alloy ingot for permanent magnet, alloy powder and method for producing bonded magnet

Info

Publication number: JP3710837B2
Application number: JP30153294A
Authority: JP
Inventors: 尚池上; 浩之冨澤; 誠一細川; 稔上原; 哲広沢; 俊郎富田
Original assignee: Neomax Co Ltd
Current assignee: Hitachi Metals Ltd
Priority date: 1994-04-07
Filing date: 1994-11-09
Publication date: 2005-10-26
Anticipated expiration: 2020-10-26
Also published as: JPH07331394A

Description

【０００１】
【産業上の利用分野】
この発明は、各種モーター、アクチュエーター等に用いることが可能な高い磁化と優れた磁気的異方性そして高保磁力を有するＲ（希土類元素）−Ｔ（鉄属元素）−（Ｍ）−Ｂ系永久磁石用希土類合金鋳塊および合金粉末、並びにボンド磁石の製造方法に係り、鋳塊の断面寸法比を特定して正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物が８５％以上で、低濃度のＲ、Ｂ領域がない鋳塊を得て、これを特定雰囲気で焼鈍し、特定の昇温条件、雰囲気条件でＨ₂ガス中で水素化処理し、さらに特定条件の脱Ｈ₂処理後に冷却することにより、結晶粒径が微細かつ磁気異方性を有する希土類合金粉末を得て、磁気特性の優れたボンド磁石を製造する永久磁石用希土類合金鋳塊および合金粉末並びにボンド磁石の製造方法に関する。
【０００２】
【従来の技術】
Ｒ−Ｔ−（Ｍ）−Ｂ系永久磁石用合金の鋳造方法として、例えば特開平２−２５１３５９号公報に、溶湯を鋳型に鋳込む際に、内面に離型剤を塗布した鋳型に１２００〜１７００℃の溶湯を鋳込み、冷却時には包晶温度前後で冷却速度を変化させ、鋳造組成をほとんど柱状晶とする方法が提案されている。
上記鋳造方法で溶製された鋳塊は、鋳造組織がほとんど一方向に配向しているため、磁気的異方性の優れた鋳造磁石ができるとされている。
【０００３】
また、永久磁石用希土類合金粉末の水素化処理法による製造方法として、Ｒ−Ｔ−（Ｍ）−Ｂ系原料合金インゴットまたは粉末を、Ｈ₂ガス雰囲気またはＨ₂ガスと不活性ガスの混合雰囲気中で温度５００℃〜１０００℃に保持して上記合金のインゴットまたは粉末にＨ₂を吸蔵させた後、Ｈ₂ガス圧力１３Ｐａ（１×１０^-1Ｔｏｒｒ）以下の真空雰囲気またはＨ₂ガス分圧１３Ｐａ（１×１０^-1Ｔｏｒｒ）以下の不活性ガス雰囲気になるまで温度５００℃〜１０００℃で脱Ｈ₂処理し、次いで冷却する水素化処理方法が提案（特開平１−１３２１０６号公報）されている。
【０００４】
前記水素化処理法により製造されたＲ−Ｔ−（Ｍ）−Ｂ系合金磁石粉末は大きな保磁力と磁気異方性を有する。これは上記処理によって非常に微細な再結晶粒径、実質的には０．１μｍ〜１μｍの平均再結晶粒径を持つ組織となり、磁気的には正方晶Ｎｄ₂Ｆｅ₁₄Ｂ系化合物の単磁区臨界粒径に近い結晶粒径となっており、かつこれらの極微細結晶がある程度結晶方位を揃えて再結晶しているためである。
【０００５】
【発明が解決しようとする課題】
このような結晶組織がボンド磁石用原料に要求される４００μｍ以下程度の粉末粒子内で単一の配向方向を有するためには、原料の鋳塊組織の異方性が高いことおよび成分の偏析が少ないことが要求され、そのためには柱状晶が発達したマクロ組織を得、かつ結晶組織が十分大きいという、鋳造では相反する条件を同時に満たすようにする必要がある。
【０００６】
ところが、上記Ｒ−Ｔ−（Ｍ）−Ｂ系永久磁石用希土類合金鋳塊の鋳造において、組成によって鋳造組織が大きく変化するため、鋳造条件を変化させても、あらゆる組成で異方性向上のためにすべての組織を柱状晶とすることは、一般に困難である。
特に本系合金鋳塊においては、目的の正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物相の組成に過剰なＲとＢを加えることで柱状晶が成長しやすくなるが、合金の磁化を向上させるためには、目的の正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物相以外の相が存在することは好ましくなく、矛盾するこの２つの条件を満たすことは非常に困難であった。
【０００７】
この発明は、Ｒ−Ｔ−（Ｍ）−Ｂ系永久磁石用希土類合金鋳塊の鋳造において、鋳塊のすべての組織部位を柱状晶と同等な異方化度をもたせることができる鋳塊の製造方法の提供、また、上記鋳塊より結晶粒径が微細で磁気異方性を有する希土類合金粉末を取扱い容易にかつ効率よく製造できる希土類合金粉末の製造方法の提供、さらに、かかる希土類合金粉末より磁気特性の優れたボンド磁石を製造するボンド磁石の製造方法の提供を目的としている。
【0008】
【課題を解決するための手段】
発明者らは、この発明を完成するに当たって、原料鋳塊の鋳造組織および組成と鋳造条件の関係を詳細に検討した結果、以下のことを明らかにした。
原料鋳塊の磁化向上、つまり鋳塊中の正方晶Nd₂Fe₁₄B型化合物相の含有率を向上させるために過剰なRとBを減少させると、得られる鋳造組織において柱状晶の割合が減少し、主にチル晶、柱状晶、等軸晶から構成されるようになる。
しかし、鋳塊サイズを限定して冷却条件を管理すると、等軸晶部分が、各々一定の成長方向に成長したデンドライト晶の集合組織となり、これによって等軸晶部でも結晶方位が同一な領域が広くなり、等軸晶部でも柱状晶部と変わりない異方性を持つ合金粉末が得られる。
さらに、このような過剰なNdとBが少ない組成ではインゴットの厚み方向に沿って鋳造組織を4種類以上に分類でき、それらはチル晶、柱状晶、等軸晶および中央部付近の鉄過剰組織である。
特に、鋳塊の厚み方向の中央部付近では合金中のNdとBの組成が大きく減少していて、正方晶Nd₂Fe₁₄B型化合物相の存在比が大きく低下し、長時間の焼鈍によっても均質化されないため、高特性の磁性粉末を得ることができない。
そこで、発明者らは、この中央部付近のNdとBが少ない鉄過剰組織が生成しない方法を種々に検討し、鋳塊の鋳造条件のうち、鋳込み方向に平行な鋳塊断面の熱流方向となる厚み方向と幅方向の寸法とその比率によって鋳造組織の制御が可能なことを知見し、この発明を完成した。
【０００９】
すなわち、この発明は、
R:11.5〜12.5at%(R:Yを含む希土類元素の少なくとも1種で、かつPrまたはNdの1種または2種をRのうち50at%以上含有)、T:79〜83at%(T:FeまたはFeの一部を50at%以下のCoで置換)、M:0.01〜2at%(M:Ga、Zr、Nb、Hf、Taのうち1種または2種以上)、B:5.5〜6.5at%からなり、鋳込み方向に平行な鋳塊の断面において、断面寸法が、幅L:30mm以上、鋳造時の熱流方向となる厚みd:10mm〜35mmで、寸法比L/dが3.0以上の関係を有し、該厚み方向で外側から中央に向かって、主にチル晶、柱状晶、等軸晶から構成され、該等軸晶部分が、各々一定の成長方向に成長したデンドライト晶の集合組織からなり、かつ正方晶Nd₂Fe₁₄B型化合物が該合金中に85%以上存在し、鋳塊中にRが11.0%及びBが5.0%を下回る領域が体積比で20%以上存在しないことを特徴とする永久磁石用希土類合金鋳塊である。
【００１０】
また、この発明は、
上記構成の永久磁石用希土類合金鋳塊を、１１２０℃〜１１６０℃で０．５時間〜１００時間、不活性雰囲気もしくは真空中で焼鈍後、平均粒度３０μｍ〜５０００μｍに粉砕し、
Ｈ₂雰囲気中の昇温過程において、６００℃〜７５０℃の温度域を昇温速度５℃／ｍｉｎ〜２００℃／ｍｉｎで昇温し、
１０ｋＰａ〜１０００ｋＰａのＨ₂ガス中で７５０℃〜９００℃に１５分〜８時間加熱保持し、組織をＲ水素化物、Ｔ−Ｂ化合物、Ｔ相、Ｒ₂Ｔ₁₄Ｂ化合物の少なくとも４相の混合組織とする水素化不均化工程の後、
さらにＡｒガスまたはＨｅガスによる絶対圧１００Ｐａ〜５０ｋＰａの減圧気流中もしくは真空排気によって炉内の水素分圧を１０ｋＰａ以下に保持しながら、７００℃〜９００℃、５分〜８時間加熱保持する脱Ｈ₂処理し、
その後冷却して平均結晶粒径が０．０５μｍ〜１μｍで、磁気的かつ結晶配向的に異方性を有することを特徴とする永久磁石用希土類合金粉末の製造方法である。
【００１１】
また、この発明は、
上記構成の製造方法で得た永久磁石用希土類合金粉末を、平均粒度２０μｍ〜４００μｍに粉砕し、この粉砕粉末に樹脂または低融点金属を混合し成形固化することを特徴とするボンド磁石の製造方法である。
【００１２】
組成の限定理由
この発明に使用する原料合金に用いるＲすなわち希土類元素は、Ｙ、Ｌａ、Ｃｅ、Ｐｒ、Ｎｄ、Ｓｍ、Ｇｄ、Ｔｂ、Ｄｙ、Ｈｏ、Ｅｒ、Ｔｍ、Ｌｕが包括され、このうち少なくとも１種以上で、Ｐｒ、Ｎｄのうち少なくとも１種または２種をＲのうち５０ａｔ％以上含有し、さらにＲの全てがＰｒ、Ｎｄのうち１種または２種の場合がある。
Ｒの５０ａｔ％以上をＰｒ、Ｎｄのうち少なくとも１種以上とするのは、５０ａｔ％未満では充分な磁化が得られないためである。
Ｒは、１１．５ａｔ％未満ではＴ相が鋳塊の中央部に大量に析出するために保磁力および角形性が低下し、また１２．５ａｔ％を越えると、目的とする正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物以外に、Ｒリッチの第２相が多く析出し、この第２相が多すぎると合金の磁化を低下させる。従ってＲの範囲は１１．５〜１２．５ａｔ％とする。好ましい範囲は１１．８〜１２．３ａｔ％である。
【００１３】
Ｔは鉄属元素であって、Ｆｅ、Ｃｏを包含する。Ｔが７９ａｔ％未満では低保磁力、低磁化の第２相が晶出して磁気的特性が低下し、また、８３ａｔ％を超えるとＴ相の晶出により保磁力、角型性が低下するため、Ｔの範囲は７９〜８３ａｔ％とする。
また、Ｆｅのみでも必要な磁気的性質は得られるが、Ｃｏの適量の添加は、キュリー温度の向上に有用であり、Ｃｏは必要に応じて添加できる。ＦｅとＣｏの原子比においてＦｅが５０％以下となるとＮｄ₂Ｆｅ₁₄Ｂ型化合物の飽和磁化そのものの減少量が大きくなってしまうため、Ｔのうち原子比でＦｅを５０％以上とした。Ｔの好ましい範囲は８０〜８２ａｔ％である。
【００１４】
添加元素Ｍの効果は、水素化時に母相の分解反応を完全に終了させずに、母相すなわちＲ₂Ｔ₁₄Ｂ相を安定化して故意に残存させるのに有効な元素が望まれる。特に顕著な効果を持つものとして、Ｇａ、Ｚｒ、Ｈｆ、Ｔａ、Ｎｂがある。
添加量は、０．０１ａｔ％未満では異方性が低下し、また２．０ａｔ％を越えると強磁性でない第２相が析出して磁化を低下させることから、Ｍは０．０１〜２．０ａｔ％とした。Ｍの好ましい範囲は０．５〜１．５ａｔ％である。
【００１５】
Ｂは、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型結晶構造を安定して析出させるためには必須の元素である。添加量は５．５ａｔ％以下ではＲ₂Ｔ₁₇相が析出して保磁力を低下させ、減磁曲線の角型性が著しく損なわれる。また、６．５ａｔ％を超えて添加した場合は、磁化の小さい第２相が析出して粉末の磁化を低下させる。従って、Ｂは、５．５〜６．５ａｔ％とした。好ましい範囲は５．８〜６．３ａｔ％である。
【００１６】
製造条件の限定理由
鋳造方法の限定理由
一般に鋳造組織は鋳塊の組成および鋳造条件によって決まり、この発明においては鋳塊形状によってさらにこれを制御することを特徴としている。
この発明において、鋳塊断面の寸法、厚み：ｄを１０〜３５ｍｍとしたのは、溶湯の入った坩堝から鋳型に鋳込む際に、鋳型厚みが１０ｍｍ未満では実際の工程においてハンドリングの面で困難がともなったり、また生産性が悪いためコスト高になるので好ましくなく、また、３５ｍｍを越えると、鋳塊幅：Ｌが実質的に無限大であっても鋳塊の厚み方向における中央部にＴの初晶が多く晶出するＮｄとＢの成分が少ない組織となり、この領域では正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物相の存在比が低下するために好ましくない。よって鋳塊厚み：ｄを１０〜３５ｍｍとする。
また、鋳塊断面幅：Ｌと寸法比Ｌ／ｄは、ｄが例えば１０ｍｍあってもＬが３０ｍｍ未満、つまりＬ／ｄ＜３．０では鋳塊を鋳造する際、厚み方向以外からの熱流方向に影響される領域が多くなり、鋳塊の異方性の発現要因であるＲ₂Ｆｅ₁₄Ｂ相の結晶方位が同方向を向いた領域が小さくなり、異方性が低下する。よって、Ｌを３０ｍｍ以上で、かつＬ／ｄが３．０以上であることが必要である。
【００１７】
この発明において、鋳型の材質構造は、この発明の対象組成の合金が鋳造できる金属、耐火物などであれば特に限定しない。鋳造条件、特に必要とされる冷却速度に応じて選択すればよく、鋳塊厚みがそれほど厚くなく、寸法比Ｌ／ｄが充分に大きい場合には空冷鉄製鋳型で充分であり、必要によっては、銅製の水冷鋳型を用いてもよい。
この発明で使用する鋳型の厚みは、鋳塊厚みに対して必要とされる厚さにすればよい。具体的には鋳塊厚みのおよそ１／２〜２倍の範囲が好ましい。実用上は５〜４０ｍｍの範囲で選択すれば充分である。
【００１８】
この発明において、原料合金中の正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の含有量は、該化合物が８５ｖｏｌ％未満であると、磁気特性が低下する。より具体的には、混在する第２相がＴの初晶の場合は保磁力を低下させ、Ｒリッチ相やＢリッチ相の場合には磁化が低下するため、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の存在比を８５ｖｏｌ％以上とした。
【００１９】
また、原料鋳塊中でＲとＢの成分が、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の化学量論比を大きく下回る、つまり、Ｒ＜１１．０ａｔ％およびＢ＜５．０ａｔ％となる組織が存在すると、焼鈍後もＴの初晶などの保磁力や角型性を低下させる相が多く残存する。従ってＲとＢの成分がそれぞれ１１．０ａｔ％および５．０ａｔ％を下回る組織が２０％以上存在しないことが必要である。
【００２０】
粉末の製造方法の限定理由
この発明の鋳塊の焼鈍条件において、焼鈍温度を１１２０℃〜１１６０℃としたのは、１１２０℃未満では固相反応が主体となり、拡散速度が十分速くないためにマクロな偏析は解消されず、粒成長や元素の拡散による均質化が充分でなく、粉砕した時の粉末全体での異方性が低くなり、また、１１６０℃を越えると焼鈍中に液相が大量に発生して焼鈍に用いた容器と反応したり、設備が大がかりになるため好ましくない。よって焼鈍温度は１１２０℃〜１１６０℃とする。好ましくは１１３０℃〜１１５５℃である。
焼鈍時間を０．５時間〜１００時間としたのは、０．５時間未満では焼鈍時の粒成長や元素の拡散による均質化が充分でなく、また、１００時間を越えて行っても、１００時間焼鈍した場合に比べて効果面で顕著に変化しない上に、長時間の焼鈍は実質的なコスト高になる。よって焼鈍時間を０．５時間〜１００時間とする。好ましくは４時間〜２４時間である。
【００２１】
この発明の出発原料の粗粉砕方法は、従来の機械的粉砕法の他、Ｈ₂吸蔵による自然崩壊、いわゆる水素粉砕などいずれの方法でもよい。
粗粉砕粉の平均粒度を３０μｍ〜５０００μｍに限定したのは、平均粒度が３０μｍ未満では粉末の酸化による磁性劣化の恐れがあり、また、５０００μｍを越えるとＲ水素化物相、Ｔ相、Ｔ−Ｂ相などに相分離する水素化、不均化反応の進行時間で局所的な差異が生じ、大きな異方性を持たせるのが困難となるからである。
【００２２】
水素化処理法とは、所要粒度の粗粉砕粉が外観上その大きさを変化させることなく、極微細結晶組織の集合体が得られることを特徴とする。
すなわち、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物に対し、高温、実際上は６００℃〜９００℃の温度範囲でＨ₂ガスと反応させると、Ｒ水素化物相、Ｔ相、Ｔ−Ｂ相などに相分離し、さらに同温度域でＨ₂ガスを脱Ｈ₂処理により除去すると、再度正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の再結晶組織が得られる。
しかしながら、現実には、水素化不均化処理条件によって分解生成物の結晶粒径、反応の度合いが異なり、水素化不均化状態の金属組織は、水素化温度７５０℃未満と７５０℃以上で明らかに異なる。この金属組織上の違いが、脱水素処理を行った後の粉末の磁気的性質、特に磁気異方性に大きく影響する。
また、水素化処理する前の鋳塊の状態、特に粒度が脱水素処理を行った後の磁粉の磁気的性質特に磁気異方性に大きく影響する。
さらに、脱水素処理条件によって、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の再結晶状態が大きく影響を受け、水素処理法によって作製した磁紛の磁気的性質、特に保磁力に大きく影響する。
【００２３】
この発明において、Ｈ₂ガス中での昇温速度は、５℃／ｍｉｎ未満であると、昇温過程で６００℃〜７５０℃の温度域を、分解反応が進行しながら通過するために、完全に分解して母相すなわちＲ₂Ｔ₁₄Ｂ相が残存せず、脱水素処理後の磁気的及び結晶方位的異方性がほとんど失われてしまう。
また、処理条件によっては、大きな反応熱のために局部的に最適処理温度範囲を越える場合があり、そのために実用的な保磁力が得られない場合がある。昇温速度を５℃／ｍｉｎ以上にすれば、６００℃〜７５０℃の領域で反応が充分に進行せず、母相を残存したまま７５０℃〜９００℃の水素化温度域に達するため、脱水素処理後に磁気的および結晶方位的に大きな異方性を持った粉末を得ることができる。従って、昇温速度は、７５０℃以下の温度域において、５℃／ｍｉｎ以上とする必要がある。
また、２００℃／ｍｉｎを越える昇温速度は赤外線加熱炉等を用いても実質的に実現困難であり、また可能であっても設備費が過大となり好ましくない。よって、昇温速度を５℃〜２００℃／ｍｉｎとする。
【００２４】
この発明において、水素化工程におけるＨ₂ガス中での保持に際し、Ｈ₂ガス圧力が１０ｋＰａ未満では、前述の分解反応が充分に進行せず、また１０００ｋＰａを越えると処理設備が大きくなりすぎ、工業的にコスト面、また安全面で好ましくないため、圧力範囲を１０ｋＰａ〜１０００ｋＰａとした。さらに好ましくは５０ｋＰａ〜１５０ｋＰａである。
【００２５】
水素化工程におけるＨ₂ガス中での加熱処理温度は、６００℃未満ではＲ水素化物相、Ｔ相、Ｔ−Ｂ相などへの分解反応が起こらず、また、６００〜７５０℃の温度範囲ではＲ水素化物の生成反応速度が速いために水素化、分解反応がほぼ完全に進行してしまい、分解生成物中に適量のＲ₂Ｔ₁₄Ｂ相が残存せず、脱水素処理後に磁気的、また結晶方位的に充分な異方性が得られない。また、９００℃を超えるとＲ水素化物相が不安定となり、かつ生成物が粒成長して正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物極微細結晶組織を得ることが困難になる。
水素化の温度範囲が７５０℃〜９００℃の領域であれば、脱水素時の再結晶反応の核となるＲ₂Ｔ₁₄Ｂ相が分散して適量残存するため、脱水素後のＲ₂Ｔ₁₄Ｂ相の結晶方位が残存Ｒ₂Ｔ₁₄Ｂ相によって決定され、結果的に再結晶組織の結晶方位が原料インゴットの結晶方位と一致し、少なくとも原料インゴットの結晶粒径の範囲内では大きな異方性を示すことになる。そのため水素化処理の温度範囲を７５０℃〜９００℃とする。
また、加熱処理保持時間については、上記の分解反応を充分に行わせるためには１５分以上必要であり、また、８時間を超えると残存Ｒ₂Ｔ₁₄Ｂ相が減少して脱水素処理後の磁気異方性が低下するため好ましくない。従って、１５分〜８時間の加熱保持とする。
【００２６】
この発明の脱Ｈ₂処理は、不活性ガス、具体的にはＡｒガスまたはＨｅガス雰囲気の減圧下もしくは真空排気で行うが、これによって粉末の周囲の実質的なＨ₂分圧はほぼ平衡水素圧、例えば８５０℃にて１ｋＰａ程度となり、脱水素反応は徐々に進行する。
不活性ガスとしてＡｒまたはＨｅに限定したのは、コスト面ではＡｒが使い良く、また、Ｈ₂ガスの置換性や温度制御性の点からはＨｅガスが優れているためである。その他の希ガスは、性能面でのメリットがない上、コスト的に問題がある。
また、一般に不活性ガスとして取り扱われるＮ₂ガスは、希土類系化合物と反応して窒化物を形成するため不適当である。
減圧気流時の雰囲気の絶対圧力が１００Ｐａ未満では、脱水素反応が急激に起こり、化学反応による温度低下が大きく、さらに、脱水素反応が急激すぎるために、冷却後の磁粉の組織に粗大な結晶粒が混在してしまい、そのために保磁力が大きく低下する。一方、雰囲気の絶対圧力が５０ｋＰａを越えると、脱水素反応に時間がかかりすぎて実用的には問題となる。そこで、雰囲気の絶対圧力は１００Ｐａ〜５０ｋＰａとした。
【００２７】
また、脱水素処理時に減圧気流中で行うのは、脱水素反応によって原料から放出されるＨ₂ガスによって、炉内圧力が上昇するのを防止するためである。実用上は一方から不活性ガスを導入しつつ、他方から真空ポンプで排気し、圧力の制御は供給口、排気口それぞれに取り付けられた流量調整弁を用いて行う。
また、真空排気で脱水素を行う場合、脱水素反応速度、すなわち、原料から放出される水素と真空排気速度で圧力を制御する。このときの圧力（水素分圧）が平衡水素圧から大きく外れると反応速度が変化し、磁粉中の組織に粗大な粒が混在したり、急激な脱水素による吸熱で原料温度が低下し、Ｒ₂Ｆｅ₁₄Ｂ相への再結晶反応が不完全となり、保磁力が大きく低下する。そこで、水素分圧は、１０ｋＰａ以下とした。
【００２８】
この発明において、脱Ｈ₂処理の温度が７００℃未満では、Ｒ水素化物相からのＨ₂の離脱が起こらないか、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ相化合物の再結晶が充分進行しない。また、９００℃を越えると正方晶Ｎｄ₂Ｆｅ₁₄Ｂ相化合物は生成するが、再結晶粒が粗大に成長し、高い保磁力が得られない。そのため、脱Ｈ₂処理の温度範囲は７００℃〜９００℃とする。
また、加熱処理保持時間は、処理設備の排気能力にもよるが、上記の再結晶反応を充分に行わせるためには少なくとも５分以上保持する必要がある。しかし、一方では、二次的な再結晶反応によって結晶が粗大化すれば保磁力の低下を招くので、できる限り短時間のほうが好ましい。そのため、５分〜８時間の加熱保持で充分である。
脱Ｈ₂処理は、原料の酸化防止の観点から、また処理設備の熱効率の観点からも、水素化処理時に引き続いて行うのがよいが、水素化処理後、一旦原料を冷却して、再び改めて脱Ｈ₂のための熱処理を行っても良い。
【００２９】
脱Ｈ₂処理後の正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の再結晶粒径は、実質的に０．０５μｍ以下の平均再結晶粒径を得ることは困難であり、またたとえ得られたとしても磁気特性上の利点がない。一方、平均再結晶粒径が１μｍを越えると、粉末の保磁力が低下するため好ましくない。そのため、平均再結晶粒径を０．０５μｍ〜１μｍとした。
【００３０】
ボンド磁石の製造方法の限定理由
この発明において、上述の製造方法による希土類合金粉末をボンド磁石用原料として粉砕する方法は従来からの機械的粉砕方法でよい。
この発明において、ボンド磁石を製造するのに用いる粉末の平均粒度を２０μｍ〜４００μｍに限定したのは、２０μｍ未満では粉末の酸化による磁気特性の劣化の恐れがあり、また、４００μｍを越えると小型磁気部品として精密成形する際に粗大すぎて好ましくないからである。
【００３１】
この発明による永久磁石合金粉末を用いて磁石化するには、以下に示す圧縮成形、射出成形、押し出し成形、圧延成形、樹脂含浸法など公知のいずれの製造方法であってもよい。
圧縮成形の場合は、磁性粉末に熱硬化性樹脂、カップリング剤、滑剤などを添加混連した後、圧縮成形して加熱樹脂を硬化して得られる。また、樹脂の代わりにＺｎ，Ａｌ等の低融点金属を用いてもよい。
射出成形、押し出し成形、圧延成形の場合は、磁性粉末に熱可塑性樹脂、カップリング剤、滑剤などを添加混連した後、射出成形、押し出し成形、圧延成形のいずれかの方法にて成形して得られる。
樹脂含浸法においては、磁性粉末を圧縮成型後、必要に応じて熱処理した後、熱硬化性樹脂を含浸させ、加熱して樹脂を硬化させて得る。また、磁性粉末を圧縮成型後、必要に応じて熱処理した後、熱可塑性樹脂を含浸させて得る。
この発明において、ボンド磁石中の磁性粉末の重量比は、前記製法により異なるが、７０〜９９．５ｗｔ％であり、残部の０．５〜３０ｗｔ％が樹脂その他である。圧縮成型の場合、磁性粉末の重量比は９５〜９９．５ｗｔ％、射出成型の場合、磁性粉末の充填率は９０〜９５ｗｔ％、樹脂含浸法の場合、磁性粉末の重量比は、９６〜９９．５ｗｔ％が好ましい。
この発明における樹脂は、熱硬化性、熱可塑性のいずれの性質を有するものも利用できるが、熱的に安定な樹脂が好ましく、例えば、ポリアミド、ポリイミド、フェノール樹脂、弗素樹脂、けい素樹脂、エポキシ樹脂などを適宜選定できる。
【００３２】
【作用】
この発明は、Ｒ−Ｔ−（Ｍ）−Ｂ系永久磁石用希土類合金鋳塊の製造に際し、鋳塊の断面寸法比を特定することにより、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物が８５％以上で低濃度のＲ、Ｂ領域が２０％以上存在しない鋳塊を得て、組織を実質的に柱状晶化するもので、例えば、等軸晶部でも柱状晶部と変わりない異方性を有し、所定の焼鈍によって均質化されて高磁気特性の磁石合金粉末を得ることができる。
また、この発明は、上記の鋳塊を特定雰囲気で焼鈍し、特定の昇温条件、雰囲気条件でＨ₂ガス中で水素化処理し、さらに特定条件の脱Ｈ₂処理後に冷却することにより、結晶粒径が微細かつ磁気異方性を有する希土類合金粉末を得て、優れた異方性と高い磁化、保磁力を有するボンド磁石を製造することが可能である。
【００３３】
【実施例】
実施例１
高周波誘導溶解法によって得られた表１に示すＮｏ．１〜８の組成の溶湯を表４に示す寸法の鉄製鋳型に鋳込むことで溶製した。
このときのＮｏ．４鋳塊の厚み方向におけるＲとＢの成分変化を図１に示す。
この鋳塊を表４に示す熱処理条件にてＡｒ雰囲気中で焼鈍して、鋳塊中の正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物の体積比を９０％以上とした。さらに、この鋳塊を、Ａｒガス雰囲気中（Ｏ₂量０．５％以下）でスタンプミルにて表４に示す平均粒度に粗粉砕した後、この粗粉砕粉を管状炉に入れ、１Ｐａ以下にまで真空排気した。
その後、純度９９．９９９９％以上のＨ₂ガスを導入しつつ、表４に示す水素化処理条件で処理を行った。こうして得た水素化原料を、引き続き表４に示す脱水素処理条件に従って脱水素処理を行った。排気にはロータリーポンプを用いた。
冷却後、原料温度が５０℃以下となったところで原料を取り出した。このときの磁石合金粉末の磁気特性を表４に示す。
【００３４】
実施例２
実施例１で得られた表４のＮｏ．１０の磁石合金粉末を、Ａｒガス雰囲気中（Ｏ₂量０．５％以下）でスタンプミルにて１５０μｍの平均粒度の粉末に粉砕した後、２．５ｗｔ％のクレゾールノボラソク樹脂を混合し、１．２ＭＡ／ｍの磁界中で０．６ＧＰａの圧力を印加して成型した。
得られた圧粉体は１６０℃のＡｒ雰囲気中で１時間硬化させ、１０ｍｍ角の立方体ボンド磁石とした。ＢＨトレーサにより測定した磁気特性を表２に示す。
【００３５】
比較例１
表１に示す実施例１と同様の組成を有する８種類の組成の溶湯を、表５に示す寸法の鉄製鋳型に鋳込んで鋳造した。
このときのＮｏ．４鋳塊の厚み方向におけるＮｄとＢの成分変化を図２に示す。
こうして得られた鋳塊を表５に示す熱処理条件にてＡｒ雰囲気中で焼鈍して、Ａｒガス雰囲気中（Ｏ₂量０．５％以下）でスタンプミルにて表５に示す平均粒度に粗粉砕した後、管状炉に入れ、１Ｐａ以下にまで真空排気した。
その後、純度９９．９９９９％以上のＨ₂ガスを導入しつつ、表５に示す処理条件で水素化処理および脱水素処理を行った。この時の磁石合金粉末の磁気特性を表５に示す。
【００３６】
比較例２
比較例１で得られたＮｏ．１０の粉末を、Ａｒガス雰囲気中（Ｏ₂量０．５％以下）でスタンプミルにて１５０μｍの平均粒度の粉末に粉砕した後、２．５ｗｔ％のクレゾールノボラック樹脂を混合し、１．２ＭＡ／ｍの磁界中で０．６ＧＰａの圧力を印加して成形した。得られた圧粉体は１６０℃のＡｒ雰囲気中で１時間硬化させ、１０ｍｍ角の立方体ボンド磁石とした。ＢＨトレーサにより測定した磁気特性を表３に示す。
【００３７】
【表１】

【００３８】
【表２】

【００３９】
【表３】

【００４０】
【表４】

【００４１】
【表５】

【００４２】
【発明の効果】
この発明は、
この発明は、Ｒ−Ｔ−（Ｍ）−Ｂ系永久磁石用希土類合金鋳塊の製造に際し、鋳塊の断面寸法比を特定することにより、正方晶Ｎｄ₂Ｆｅ₁₄Ｂ型化合物が８５％以上で低濃度のＲ、Ｂ領域がない鋳塊を得て、組織を実質的に柱状晶化するもので、例えば、等軸晶部でも柱状晶部と変わりない異方性を有し、所定の焼鈍によって均質化されて高磁気特性の磁性粉末を得ることができる。
また、この発明は、上記の鋳塊を特定雰囲気で焼鈍し、特定の昇温条件、雰囲気条件でＨ₂ガス中で水素化処理し、さらに特定条件の脱Ｈ₂処理後に冷却することにより、結晶粒径が微細かつ磁気異方性を有する希土類合金粉末を得て、優れた異方性と高い磁化、保磁力を有するボンド磁石を製造することが可能である。
【図面の簡単な説明】
【図１】実施例Ｎｏ．４鋳塊の厚み方向におけるＮｄとＢの成分（ａｔ％）変化を示すグラフである。
【図２】比較例Ｎｏ．４鋳塊の厚み方向におけるＮｄとＢの成分（ａｔ％）変化を示すグラフである。[0001]
[Industrial application fields]
The present invention is an R (rare earth element) -T (iron group element)-(M) -B permanent having high magnetization, excellent magnetic anisotropy and high coercive force that can be used for various motors, actuators and the like. The present invention relates to a rare earth alloy ingot and alloy powder for a magnet, and a method for producing a bonded magnet. ₂ Fe ₁₄ An ingot having a B-type compound of 85% or more and no low concentration of R and B regions is obtained, and this is annealed in a specific atmosphere. ₂ Hydrotreating in gas, and de-H ₂ By cooling after the treatment, a rare earth alloy powder having a fine crystal grain size and magnetic anisotropy is obtained, and a rare earth alloy ingot and alloy powder for permanent magnets for producing a bonded magnet having excellent magnetic properties, and a bonded magnet It relates to a manufacturing method.
[0002]
[Prior art]
As a casting method for an RT- (M) -B permanent magnet alloy, for example, in Japanese Patent Laid-Open No. 2-251359, when casting a molten metal into a mold, a mold having a release agent applied to the inner surface is 1200-200. A method has been proposed in which a molten metal at 1700 ° C. is cast and the cooling rate is changed around the peritectic temperature during cooling so that the casting composition is almost columnar.
The ingot produced by the above casting method has a cast structure that is oriented almost in one direction, so that a cast magnet having excellent magnetic anisotropy can be produced.
[0003]
Further, as a method for producing a rare earth alloy powder for a permanent magnet by a hydrogenation method, an RT- (M) -B-based material alloy ingot or powder is used as an H ₂ Gas atmosphere or H ₂ In an atmosphere of mixed gas and inert gas, the temperature is maintained at 500 ° C. to 1000 ° C. ₂ After occlusion of H ₂ Gas pressure 13Pa (1 × 10 ^-1 Torr) The following vacuum atmosphere or H ₂ Gas partial pressure 13 Pa (1 × 10 ^-1 Torr) Degassing at a temperature of 500 ° C. to 1000 ° C. until the following inert gas atmosphere is reached ₂ A hydrotreating method for treating and then cooling has been proposed (Japanese Patent Laid-Open No. 1-132106).
[0004]
The RT- (M) -B alloy magnet powder produced by the hydrogenation method has a large coercive force and magnetic anisotropy. This results in a structure having a very fine recrystallized grain size, substantially an average recrystallized grain size of 0.1 μm to 1 μm by the above treatment, and is magnetically tetragonal Nd. ₂ Fe ₁₄ This is because the crystal grain size is close to the single domain critical grain size of the B-based compound, and these ultrafine crystals are recrystallized with a certain degree of crystal orientation.
[0005]
[Problems to be solved by the invention]
In order for such a crystal structure to have a single orientation direction in powder particles of about 400 μm or less required for a raw material for bonded magnets, the ingot structure of the raw material has high anisotropy and segregation of components. In order to achieve this, it is necessary to obtain a macrostructure in which columnar crystals are developed and to satisfy the contradictory conditions at the same time in casting that the crystal structure is sufficiently large.
[0006]
However, in casting the rare earth alloy ingot for the RT- (M) -B permanent magnet, the cast structure greatly changes depending on the composition. Therefore, even if the casting conditions are changed, the anisotropy is improved with any composition. Therefore, it is generally difficult to make all the structures into columnar crystals.
Especially in this alloy ingot, the target tetragonal Nd ₂ Fe ₁₄ Adding excessive R and B to the composition of the B-type compound phase facilitates the growth of columnar crystals, but in order to improve the magnetization of the alloy, the target tetragonal Nd ₂ Fe ₁₄ It is not preferable that a phase other than the B-type compound phase exists, and it has been very difficult to satisfy these two contradictory conditions.
[0007]
The present invention relates to an ingot that can give all the structural parts of an ingot an anisotropic degree equivalent to that of a columnar crystal in casting of a rare earth alloy ingot for an RT- (M) -B permanent magnet. Providing a production method, providing a method for producing a rare earth alloy powder capable of easily and efficiently producing a rare earth alloy powder having a crystal grain size smaller than that of the ingot and having magnetic anisotropy, and further providing the rare earth alloy powder. It aims at providing the manufacturing method of the bonded magnet which manufactures the bonded magnet which was more excellent in magnetic characteristics.
[0008]
[Means for Solving the Problems]
Upon completion of the present invention, the inventors have examined the relationship between the casting structure and composition of the raw material ingot and the casting conditions in detail, and as a result, have clarified the following.
Improvement of magnetization of raw material ingot, that is, tetragonal Nd in ingot ₂ Fe ₁₄ When excessive R and B are reduced to improve the content of the B-type compound phase, the ratio of columnar crystals in the resulting cast structure is reduced, mainly composed of chill crystals, columnar crystals, and equiaxed crystals. It becomes like this.
However, when the cooling conditions are controlled by limiting the ingot size, the equiaxed crystal part becomes a dendrite crystal structure that grows in a certain growth direction. As a result, an alloy powder having anisotropy that is the same as that of the columnar crystal part even in the equiaxed crystal part is obtained.
In addition, in such a composition with a small amount of excess Nd and B, the cast structure can be classified into four or more types along the thickness direction of the ingot, and they are chill, columnar, equiaxed, and iron-rich structure near the center. It is.
In particular, In the thickness direction of the ingot Near the center, the composition of Nd and B in the alloy is greatly reduced, and tetragonal Nd ₂ Fe ₁₄ Since the abundance ratio of the B-type compound phase is greatly reduced and is not homogenized even by annealing for a long time, it is not possible to obtain a magnetic powder with high characteristics.
Therefore, the inventors examined various methods of producing an iron excess structure with little Nd and B near the center, and among the casting conditions of the ingot, Parallel to casting direction It was found that the cast structure can be controlled by the thickness direction and width direction dimensions and the ratio of the heat flow direction of the ingot cross section, and the present invention was completed.
[0009]
That is, this invention
R: 11.5 to 12.5 at% (containing at least one rare earth element including R: Y and containing one or two of Pr or Nd of 50 at% or more of R), T: 79 to 83 at% (T: Fe or a part of Fe is replaced with Co of 50 at% or less), M: 0.01 to 2 at% (M: Ga, Zr, Nb, Hf, Ta or more of Ta), B: 5.5 to 6.5 at % Parallel to casting direction Ingot In cross section Cross-sectional dimension, width L: 30mm or more, Casting The thickness d in the heat flow direction is 10 mm to 35 mm, and the dimension ratio L / d has a relationship of 3.0 or more. From the outer side to the center in the thickness direction, mainly composed of chill crystals, columnar crystals, and equiaxed crystals, and the equiaxed crystal portions are each composed of a dendrite crystal structure grown in a certain growth direction, and Tetragonal Nd ₂ Fe ₁₄ B-type compound is present in the alloy in an amount of 85% or more, and in the ingot, a region where R is 11.0% and B is less than 5.0% is not present in a volume ratio of 20% or more. It is a lump.
[0010]
The present invention also provides
The rare earth alloy ingot for permanent magnet having the above structure is pulverized to an average particle size of 30 μm to 5000 μm after annealing in an inert atmosphere or vacuum at 1120 ° C. to 1160 ° C. for 0.5 hour to 100 hours,
H ₂ In the temperature rising process in the atmosphere, the temperature is raised from 600 ° C. to 750 ° C. at a temperature rising rate of 5 ° C./min to 200 ° C./min,
H of 10 kPa to 1000 kPa ₂ Heat and hold at 750 ° C. to 900 ° C. for 15 minutes to 8 hours in gas, and the structure is R hydride, TB compound, T phase, R ₂ T ₁₄ After the hydrogenation disproportionation step to form a mixed structure of at least four phases of B compound,
Further, de-H is maintained by heating at 700 ° C. to 900 ° C. for 5 minutes to 8 hours while maintaining the hydrogen partial pressure in the furnace at 10 kPa or less in a reduced pressure air flow of 100 Pa to 50 kPa absolute pressure by Ar gas or He gas or by vacuum exhaust. ₂ Process,
Then, it is cooled, and the average crystal grain size is 0.05 μm to 1 μm, and the method is a method for producing a rare earth alloy powder for permanent magnets having magnetic and crystal orientation anisotropy.
[0011]
The present invention also provides
A rare earth alloy powder for a permanent magnet obtained by the above-described production method is pulverized to an average particle size of 20 μm to 400 μm, and the pulverized powder is mixed with a resin or a low-melting-point metal to be molded and solidified. It is.
[0012]
Reason for limitation of composition
R, that is, a rare earth element used in the raw material alloy used in the present invention includes Y, La, Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er, Tm, and Lu, and at least one of them is included. In some cases, at least one or two of Pr and Nd are contained in 50 at% or more of R, and all of R may be one or two of Pr and Nd.
The reason why at least 50 at% of R is at least one of Pr and Nd is that sufficient magnetization cannot be obtained at less than 50 at%.
If R is less than 11.5 at%, a large amount of T phase precipitates in the center of the ingot, so that the coercive force and squareness are reduced. If it exceeds 12.5 at%, the desired tetragonal Nd ₂ Fe ₁₄ In addition to the B-type compound, many R-rich second phases are precipitated, and if there are too many second phases, the magnetization of the alloy is lowered. Therefore, the range of R is 11.5 to 12.5 at%. A preferred range is 11.8 to 12.3 at%.
[0013]
T is an iron group element and includes Fe and Co. If T is less than 79 at%, the second phase with low coercive force and low magnetization will be crystallized and the magnetic properties will deteriorate, and if it exceeds 83 at%, coercive force and squareness will decrease due to T phase crystallization. , T ranges from 79 to 83 at%.
Moreover, necessary magnetic properties can be obtained with Fe alone, but addition of an appropriate amount of Co is useful for improving the Curie temperature, and Co can be added as necessary. When Fe is 50% or less in the atomic ratio of Fe and Co, Nd ₂ Fe ₁₄ Since the amount of decrease in the saturation magnetization itself of the B-type compound becomes large, the atomic ratio of Fe is set to 50% or more. A preferable range of T is 80 to 82 at%.
[0014]
The effect of the additive element M is that the decomposition of the matrix phase is not completely terminated during hydrogenation, and the matrix phase, that is, R ₂ T ₁₄ An element effective for stabilizing the B phase and intentionally remaining is desired. Ga, Zr, Hf, Ta, and Nb are particularly effective.
If the added amount is less than 0.01 at%, the anisotropy decreases, and if it exceeds 2.0 at%, a non-ferromagnetic second phase precipitates and decreases the magnetization. It was set to 0 at%. A preferable range of M is 0.5 to 1.5 at%.
[0015]
B is tetragonal Nd ₂ Fe ₁₄ It is an essential element for stably depositing the B-type crystal structure. Addition amount is less than 5.5at%, R ₂ T ₁₇ A phase precipitates to reduce the coercive force, and the squareness of the demagnetization curve is significantly impaired. Moreover, when adding over 6.5 at%, the 2nd phase with small magnetization precipitates and the magnetization of powder is reduced. Therefore, B is set to 5.5 to 6.5 at%. A preferred range is 5.8 to 6.3 at%.
[0016]
Reasons for limiting production conditions
Reason for limitation of casting method
In general, the cast structure is determined by the composition of the ingot and the casting conditions, and the present invention is characterized in that this is further controlled by the shape of the ingot.
In this invention, the dimension and thickness of the ingot cross-section: d is set to 10 to 35 mm. When casting from a crucible containing molten metal into a mold, if the mold thickness is less than 10 mm, it is difficult in handling in the actual process. In addition, it is not preferable because the cost is high because the productivity is poor, and when it exceeds 35 mm, the ingot width: L is substantially infinite. In this region, tetragonal Nd and Bd have a structure with few Nd and B components. ₂ Fe ₁₄ This is not preferable because the abundance ratio of the B-type compound phase is lowered. Therefore, ingot thickness: d shall be 10-35 mm.
The ingot cross-sectional width: L and the dimensional ratio L / d are such that even if d is 10 mm, for example, L is less than 30 mm, that is, when L / d <3.0, The area affected by the direction increases, and R is a factor in manifesting the anisotropy of the ingot. ₂ Fe ₁₄ The region in which the crystal orientation of the B phase is directed in the same direction is reduced, and the anisotropy is reduced. Therefore, it is necessary that L is 30 mm or more and L / d is 3.0 or more.
[0017]
In the present invention, the material structure of the mold is not particularly limited as long as it is a metal, a refractory, or the like that can cast an alloy having the target composition of the present invention. What is necessary is just to select according to casting conditions, especially the cooling rate required, and when the ingot thickness is not so thick and the dimensional ratio L / d is sufficiently large, an air-cooled iron mold is sufficient. A copper water-cooled mold may be used.
The thickness of the mold used in the present invention may be a thickness required for the ingot thickness. Specifically, a range of about 1/2 to 2 times the ingot thickness is preferable. Practically, it is sufficient to select in the range of 5 to 40 mm.
[0018]
In this invention, tetragonal Nd in the raw material alloy ₂ Fe ₁₄ When the content of the B-type compound is less than 85 vol%, the magnetic properties are deteriorated. More specifically, when the coexisting second phase is a primary crystal of T, the coercive force is reduced, and when it is an R-rich phase or a B-rich phase, the magnetization is reduced. ₂ Fe ₁₄ The abundance ratio of the B-type compound was 85 vol% or more.
[0019]
In the raw material ingot, the components of R and B are tetragonal Nd ₂ Fe ₁₄ If there is a structure that greatly falls below the stoichiometric ratio of the B-type compound, that is, R <11.0 at% and B <5.0 at%, the coercive force and squareness of the primary crystal of T will be reduced even after annealing. Many phases to decrease remain. Accordingly, it is necessary that there are no more than 20% of the structures having R and B components below 11.0 at% and 5.0 at%, respectively.
[0020]
Reasons for limiting powder production methods
In the ingot annealing conditions of this invention, the annealing temperature was set to 1120 ° C. to 1160 ° C., and the solid phase reaction was the main at less than 1120 ° C., and the macroscopic segregation was not eliminated because the diffusion rate was not fast enough. Homogenization due to grain growth and element diffusion is not sufficient, and the anisotropy of the whole powder when pulverized is low, and when it exceeds 1160 ° C, a large amount of liquid phase is generated during annealing and used for annealing. It is not preferable because it reacts with the existing container and the equipment becomes large. Therefore, annealing temperature shall be 1120 to 1160 degreeC. Preferably it is 1130 degreeC-1155 degreeC.
The annealing time is set to 0.5 to 100 hours. If the annealing time is less than 0.5 hours, the grain growth during annealing and the homogenization by element diffusion are not sufficient. Compared with the case of time annealing, the effect is not significantly changed, and the annealing for a long time is substantially expensive. Therefore, annealing time shall be 0.5 to 100 hours. Preferably, it is 4 hours to 24 hours.
[0021]
In addition to the conventional mechanical pulverization method, the starting material coarse pulverization method of the present invention includes H ₂ Any method such as natural decay by occlusion, so-called hydrogen pulverization may be used.
The average particle size of the coarsely pulverized powder is limited to 30 μm to 5000 μm. If the average particle size is less than 30 μm, there is a risk of magnetic deterioration due to oxidation of the powder, and if it exceeds 5000 μm, the R hydride phase, T phase, TB This is because local differences occur in the progress time of the hydrogenation and disproportionation reactions that cause phase separation into phases, and it becomes difficult to have large anisotropy.
[0022]
The hydrotreating method is characterized in that a coarsely pulverized powder having a required particle size can obtain an aggregate of an ultrafine crystal structure without changing its size in appearance.
That is, tetragonal Nd ₂ Fe ₁₄ Compared to the B-type compound, it is H at a high temperature, actually in the temperature range of 600 ° C to 900 ° C. ₂ When reacted with gas, it is phase-separated into R hydride phase, T phase, TB phase, etc. ₂ Degas H ₂ Once removed by treatment, again tetragonal Nd ₂ Fe ₁₄ A recrystallized structure of the B-type compound is obtained.
However, in reality, the crystal grain size of the decomposition product and the degree of reaction differ depending on the hydrogenation disproportionation treatment conditions, and the metal structure in the hydrogenation disproportionation state has a hydrogenation temperature of less than 750 ° C. and more than 750 ° C. Obviously different. This difference in the metal structure greatly affects the magnetic properties of the powder after the dehydrogenation treatment, particularly the magnetic anisotropy.
In addition, the state of the ingot before the hydrogenation treatment, particularly the particle size, greatly affects the magnetic properties, particularly the magnetic anisotropy, of the magnetic powder after the dehydrogenation treatment.
Furthermore, depending on the dehydrogenation treatment conditions, tetragonal Nd ₂ Fe ₁₄ The recrystallized state of the B-type compound is greatly affected, and greatly affects the magnetic properties of the magnetic powder produced by the hydrogen treatment method, particularly the coercive force.
[0023]
In this invention, H ₂ If the temperature rise rate in the gas is less than 5 ° C./min, it passes through the temperature range of 600 ° C. to 750 ° C. during the temperature rise process while the decomposition reaction proceeds, so that it is completely decomposed and the parent phase Ie R ₂ T ₁₄ The B phase does not remain, and the magnetic and crystal orientation anisotropy after the dehydrogenation process is almost lost.
Further, depending on the processing conditions, there may be a case where the optimum processing temperature range is locally exceeded due to a large reaction heat, so that a practical coercive force may not be obtained. If the rate of temperature increase is 5 ° C./min or more, the reaction does not proceed sufficiently in the region of 600 ° C. to 750 ° C., and the hydrogenation temperature region of 750 ° C. to 900 ° C. is reached with the parent phase remaining. After elementary treatment, a powder having large anisotropy in magnetic and crystal orientation can be obtained. Therefore, the temperature increase rate needs to be 5 ° C./min or more in a temperature range of 750 ° C. or less.
Moreover, even if it uses an infrared heating furnace etc., the temperature increase rate over 200 degrees C / min is substantially difficult to implement | achieve, and even if it is possible, an installation cost becomes excessive and is unpreferable. Therefore, the rate of temperature rise is set to 5 ° C. to 200 ° C./min.
[0024]
In this invention, H in the hydrogenation process ₂ When holding in gas, H ₂ If the gas pressure is less than 10 kPa, the above-described decomposition reaction does not proceed sufficiently, and if it exceeds 1000 kPa, the treatment facility becomes too large, which is not preferable from the viewpoint of cost and safety industrially. Therefore, the pressure range is from 10 kPa to 1000 kPa. It was. More preferably, it is 50 kPa-150 kPa.
[0025]
H in the hydrogenation process ₂ When the heat treatment temperature in the gas is less than 600 ° C., decomposition reaction into R hydride phase, T phase, TB phase, etc. does not occur, and in the temperature range of 600 to 750 ° C., R hydride formation reaction Since the speed is high, the hydrogenation and decomposition reaction proceeds almost completely, and an appropriate amount of R is contained in the decomposition product. ₂ T ₁₄ The B phase does not remain, and sufficient anisotropy cannot be obtained magnetically and crystallographically after the dehydrogenation treatment. On the other hand, when the temperature exceeds 900 ° C., the R hydride phase becomes unstable, and the product grows to form tetragonal Nd. ₂ Fe ₁₄ It becomes difficult to obtain a B-type compound ultrafine crystal structure.
If the temperature range of hydrogenation is in the range of 750 ° C. to 900 ° C., R which is the nucleus of the recrystallization reaction during dehydrogenation ₂ T ₁₄ Since B phase is dispersed and remains in an appropriate amount, R after dehydrogenation ₂ T ₁₄ B phase crystal orientation remains R ₂ T ₁₄ As a result, the crystal orientation of the recrystallized structure coincides with the crystal orientation of the raw material ingot, and shows a large anisotropy at least within the range of the crystal grain size of the raw material ingot. Therefore, the temperature range of the hydrotreatment is set to 750 ° C to 900 ° C.
In addition, the heat treatment holding time is required to be 15 minutes or longer in order to sufficiently perform the above decomposition reaction. ₂ T ₁₄ This is not preferable because the B phase is reduced and the magnetic anisotropy after the dehydrogenation treatment is lowered. Accordingly, the heating and holding is performed for 15 minutes to 8 hours.
[0026]
De-H of this invention ₂ The treatment is carried out under reduced pressure or evacuation of an inert gas, specifically Ar gas or He gas atmosphere, so that a substantial H around the powder is obtained. ₂ The partial pressure becomes approximately an equilibrium hydrogen pressure, for example, about 1 kPa at 850 ° C., and the dehydrogenation reaction proceeds gradually.
The reason why the inert gas is limited to Ar or He is that Ar is easy to use in terms of cost. ₂ This is because He gas is superior in terms of gas substitutability and temperature controllability. Other noble gases have no performance advantage and are costly.
In general, N treated as an inert gas ₂ Gases are unsuitable because they react with rare earth compounds to form nitrides.
If the absolute pressure of the atmosphere under reduced pressure air flow is less than 100 Pa, the dehydrogenation reaction occurs abruptly, the temperature drop due to the chemical reaction is large, and the dehydrogenation reaction is too rapid, resulting in coarse crystals in the magnetic powder structure after cooling. The grains are mixed, and the coercive force is greatly reduced. On the other hand, if the absolute pressure of the atmosphere exceeds 50 kPa, it takes a long time for the dehydrogenation reaction, which is practically problematic. Therefore, the absolute pressure of the atmosphere was set to 100 Pa to 50 kPa.
[0027]
In addition, what is performed in a reduced-pressure air flow during the dehydrogenation process is the H released from the raw material by the dehydrogenation reaction. ₂ This is for preventing the furnace pressure from being increased by the gas. In practice, an inert gas is introduced from one side and exhausted by a vacuum pump from the other side, and the pressure is controlled using a flow rate adjusting valve attached to each of the supply port and the exhaust port.
Further, when dehydrogenation is performed by evacuation, the pressure is controlled by the dehydrogenation reaction rate, that is, the hydrogen released from the raw material and the evacuation rate. When the pressure (partial hydrogen pressure) at this time greatly deviates from the equilibrium hydrogen pressure, the reaction rate changes, coarse particles are mixed in the structure of the magnetic powder, or the raw material temperature decreases due to endotherm due to rapid dehydrogenation. ₂ Fe ₁₄ The recrystallization reaction into the B phase becomes incomplete, and the coercive force is greatly reduced. Therefore, the hydrogen partial pressure was set to 10 kPa or less.
[0028]
In this invention, de-H ₂ When the treatment temperature is less than 700 ° C., H from the R hydride phase ₂ The separation of tetragonal Nd ₂ Fe ₁₄ Recrystallization of the B phase compound does not proceed sufficiently. When the temperature exceeds 900 ° C., tetragonal Nd ₂ Fe ₁₄ Although the B phase compound is formed, the recrystallized grains grow coarsely and a high coercive force cannot be obtained. Therefore, de-H ₂ The temperature range of processing shall be 700 to 900 degreeC.
In addition, the heat treatment holding time depends on the exhaust capacity of the processing equipment, but it is necessary to hold the heat treatment for at least 5 minutes in order to sufficiently perform the recrystallization reaction. However, on the other hand, if the crystal is coarsened by a secondary recrystallization reaction, the coercive force is lowered, so that a shorter time is preferable. Therefore, heating and holding for 5 minutes to 8 hours is sufficient.
De-H ₂ From the viewpoint of preventing oxidation of the raw material and from the viewpoint of thermal efficiency of the processing equipment, the treatment is preferably carried out continuously during the hydrogenation treatment, but after the hydrogenation treatment, the raw material is once cooled and de-H again. ₂ You may perform the heat processing for.
[0029]
De-H ₂ Tetragonal Nd after treatment ₂ Fe ₁₄ It is difficult to obtain an average recrystallized grain size of 0.05 μm or less for the recrystallized grain size of the B-type compound, and even if obtained, there is no advantage in magnetic properties. On the other hand, if the average recrystallized grain size exceeds 1 μm, the coercive force of the powder is lowered, which is not preferable. Therefore, the average recrystallized grain size is set to 0.05 μm to 1 μm.
[0030]
Reason for limitation of manufacturing method of bonded magnet
In the present invention, the conventional mechanical pulverization method may be used as a method for pulverizing the rare earth alloy powder by the above-described production method as a raw material for the bond magnet.
In this invention, the average particle size of the powder used for manufacturing the bonded magnet is limited to 20 μm to 400 μm. If it is less than 20 μm, there is a risk of deterioration of magnetic properties due to oxidation of the powder. This is because it is too coarse when precision molding as a part.
[0031]
In order to magnetize the permanent magnet alloy powder according to the present invention, any of the known production methods such as compression molding, injection molding, extrusion molding, rolling molding, and resin impregnation as shown below may be used.
In the case of compression molding, the magnetic powder is obtained by adding a thermosetting resin, a coupling agent, a lubricant and the like to the magnetic powder and then compressing and curing the heated resin. Moreover, you may use low melting-point metals, such as Zn and Al, instead of resin.
In the case of injection molding, extrusion molding, and rolling molding, a thermoplastic resin, a coupling agent, a lubricant, etc. are added to and mixed with magnetic powder, and then molded by any of injection molding, extrusion molding, or rolling molding. can get.
In the resin impregnation method, the magnetic powder is compression-molded, heat-treated as necessary, impregnated with a thermosetting resin, and heated to cure the resin. Further, after compression molding, the magnetic powder is heat treated as necessary, and then impregnated with a thermoplastic resin.
In this invention, the weight ratio of the magnetic powder in the bonded magnet varies depending on the production method, but is 70 to 99.5 wt%, and the remaining 0.5 to 30 wt% is resin or the like. In the case of compression molding, the weight ratio of the magnetic powder is 95 to 99.5 wt%, in the case of injection molding, the filling ratio of the magnetic powder is 90 to 95 wt%, and in the case of the resin impregnation method, the weight ratio of the magnetic powder is 96 to 99 0.5 wt% is preferable.
As the resin in this invention, those having both thermosetting and thermoplastic properties can be used, but thermally stable resins are preferable, for example, polyamide, polyimide, phenol resin, fluorine resin, silicon resin, epoxy Resin etc. can be selected suitably.
[0032]
[Action]
In the production of a rare earth alloy ingot for an RT- (M) -B permanent magnet, the present invention specifies a tetragonal crystal Nd by specifying the cross-sectional dimension ratio of the ingot. ₂ Fe ₁₄ An ingot is obtained in which a B-type compound is 85% or more and a low concentration of R and B regions is not present in 20% or more, and the structure is substantially columnar crystallized. Magnetic alloy powder having high anisotropy and homogenized by predetermined annealing and having high magnetic properties can be obtained.
In addition, the present invention anneals the above ingot in a specific atmosphere, and H under specific temperature rise conditions and atmospheric conditions. ₂ Hydrotreating in gas, and de-H ₂ By cooling after the treatment, it is possible to obtain a rare earth alloy powder having a fine crystal grain size and magnetic anisotropy, and to produce a bonded magnet having excellent anisotropy, high magnetization and coercive force.
[0033]
【Example】
Example 1
Nos. Shown in Table 1 obtained by the high frequency induction dissolution method. It melted by casting the molten metal of the composition of 1-8 into the iron mold of the dimension shown in Table 4.
No. at this time. FIG. 1 shows changes in R and B components in the thickness direction of the four ingots.
The ingot was annealed in an Ar atmosphere under the heat treatment conditions shown in Table 4, and tetragonal Nd in the ingot was ₂ Fe ₁₄ The volume ratio of the B-type compound was 90% or more. Further, this ingot is placed in an Ar gas atmosphere (O ₂ After roughly pulverizing to an average particle size shown in Table 4 with a stamp mill, the coarsely pulverized powder was put into a tubular furnace and evacuated to 1 Pa or less.
Thereafter, H with a purity of 99.9999% or more ₂ The treatment was performed under the hydrotreating conditions shown in Table 4 while introducing the gas. The hydrogenation raw material thus obtained was subsequently dehydrogenated according to the dehydrogenation conditions shown in Table 4. A rotary pump was used for exhaust.
After cooling, the raw material was taken out when the raw material temperature became 50 ° C. or lower. Table 4 shows the magnetic properties of the magnet alloy powder at this time.
[0034]
Example 2
No. of Table 4 obtained in Example 1 was obtained. 10 magnetic alloy powders in an Ar gas atmosphere (O ₂ Crushed into a powder having an average particle size of 150 μm by a stamp mill at a quantity of 0.5% or less), and then mixed with 2.5 wt% of cresol noborazok resin and 0.6 GPa in a magnetic field of 1.2 MA / m. Molding was performed by applying pressure.
The obtained green compact was cured in an Ar atmosphere at 160 ° C. for 1 hour to obtain a 10 mm square cubic bonded magnet. Table 2 shows the magnetic properties measured by the BH tracer.
[0035]
Comparative Example 1
Eight types of molten metal having the same composition as Example 1 shown in Table 1 were cast into an iron mold having the dimensions shown in Table 5.
No. at this time. Changes in Nd and B components in the thickness direction of the four ingots are shown in FIG.
The ingot thus obtained was annealed in an Ar atmosphere under the heat treatment conditions shown in Table 5, and then in an Ar gas atmosphere (O ₂ The amount was 0.5% or less) and coarsely pulverized to an average particle size shown in Table 5 using a stamp mill, and then placed in a tubular furnace and evacuated to 1 Pa or less.
Thereafter, H with a purity of 99.9999% or more ₂ While introducing the gas, the hydrogenation treatment and the dehydrogenation treatment were performed under the treatment conditions shown in Table 5. Table 5 shows the magnetic properties of the magnet alloy powder at this time.
[0036]
Comparative Example 2
No. obtained in Comparative Example 1 10 powder in an Ar gas atmosphere (O ₂ Crushed into a powder with an average particle size of 150 μm by a stamp mill at a quantity of 0.5% or less), and then mixed with 2.5 wt% cresol novolac resin and applied a pressure of 0.6 GPa in a magnetic field of 1.2 MA / m. Applied to form. The obtained green compact was cured in an Ar atmosphere at 160 ° C. for 1 hour to obtain a 10 mm square cubic bonded magnet. Table 3 shows the magnetic characteristics measured by the BH tracer.
[0037]
[Table 1]

[0038]
[Table 2]

[0039]
[Table 3]

[0040]
[Table 4]

[0041]
[Table 5]

[0042]
【The invention's effect】
This invention
In the production of a rare earth alloy ingot for an RT- (M) -B permanent magnet, the present invention specifies a tetragonal crystal Nd by specifying the cross-sectional dimension ratio of the ingot. ₂ Fe ₁₄ An ingot that has a B-type compound of 85% or more and does not have low concentrations of R and B regions, and that has a substantially columnar crystallized structure. And can be homogenized by predetermined annealing to obtain a magnetic powder with high magnetic properties.
In addition, the present invention anneals the above ingot in a specific atmosphere, and H under specific temperature rise conditions and atmospheric conditions. ₂ Hydrotreating in gas, and de-H ₂ By cooling after the treatment, it is possible to obtain a rare earth alloy powder having a fine crystal grain size and magnetic anisotropy, and to produce a bonded magnet having excellent anisotropy, high magnetization and coercive force.
[Brief description of the drawings]
FIG. It is a graph which shows the component (at%) change of Nd and B in the thickness direction of 4 ingots.
2 is a comparative example No. 2; It is a graph which shows the component (at%) change of Nd and B in the thickness direction of 4 ingots.

Claims

R: 11.5 to 12.5 at% (containing at least one rare earth element including R: Y and containing one or two of Pr or Nd of 50 at% or more of R), T: 79 to 83 at% (T: Fe or a part of Fe is replaced with Co of 50 at% or less), M: 0.01 to 2 at% (M: Ga, Zr, Nb, Hf, Ta or more of Ta), B: 5.5 to 6.5 at In the ingot cross section parallel to the casting direction , the cross-sectional dimension is width L: 30 mm or more, the thickness d in the heat flow direction during casting is 10 mm to 35 mm, and the dimension ratio L / d is 3.0 or more. From the outer side to the center in the thickness direction, it is mainly composed of chill crystals, columnar crystals, and equiaxed crystals, and each of the equiaxed crystal parts is grown from a dendrite crystal structure grown in a certain growth direction. And tetragonal Nd ₂ Fe ₁₄ B type compound is present in the alloy at 85% or more, and the ingot has a region where R is less than 11.0% and B is less than 5.0% as a volume ratio is not more than 20%. A rare earth alloy ingot for permanent magnets.

The rare earth alloy ingot for permanent magnet according to claim 1 is annealed in an inert atmosphere or vacuum at 1120 ° C to 1160 ° C for 0.5 hour to 100 hours, and then pulverized to an average particle size of 30 µm to 5000 µm, and ascended in an H ₂ atmosphere. in elevating process, 600 ° C. to 750 was raised to a temperature range at a heating rate of 5 ℃ / min~200 ℃ / min of ° C., 15 minutes to 8 hours 750 ° C. to 900 ° C. with H ₂ gas of 10kPa~1000kPa After the hydrogenation disproportionation step, which is heated and held and the structure is a mixed structure of at least four phases of R hydride, TB compound, T phase, and R ₂ T ₁₄ B compound, the absolute pressure is 100 Pa with Ar gas or He gas. while maintaining the hydrogen partial pressure in the furnace to 10kPa or less by vacuum stream or evacuation ~50kPa, 700 ℃ ~900 ℃, de H ₂ process of heating for 5 minutes to 8 hours, the average crystal then cooled to A method for producing a rare earth alloy powder for a permanent magnet having a particle size of 0.05 μm to 1 μm and anisotropy in magnetic and crystal orientation.

A rare earth alloy powder for a permanent magnet obtained by the production method according to claim 2 is pulverized to an average particle size of 20 μm to 400 μm, and a resin or a low-melting-point metal is mixed into the pulverized powder and solidified by molding. Method.