JP2004099932A

JP2004099932A - Method for manufacturing rare-earth alloy powder for sintered magnet

Info

Publication number: JP2004099932A
Application number: JP2002260252A
Authority: JP
Inventors: Takeshi Araki; 荒木　健; Hiroyuki Teramoto; 寺本　浩行; Takanori Sone; 曽根　孝典
Original assignee: Mitsubishi Electric Corp
Current assignee: Mitsubishi Electric Corp
Priority date: 2002-09-05
Filing date: 2002-09-05
Publication date: 2004-04-02
Anticipated expiration: 2022-09-05
Also published as: JP3848906B2

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for manufacturing a rare-earth alloy powder for a sintered magnet, which reduces an amount of oxygen taken into the powder of a cast alloy, avoids a decrease of an adulterant effect and a decrease of sinterability, and prevents a lowering of residual magnetization. <P>SOLUTION: The method for manufacturing the rare-earth alloy powder for the sintered magnet comprises the step (1) of casting an alloy of R<SB>x</SB>T<SB>1-x-y</SB>B<SB>y</SB>(wherein, R is one or more elements selected among Nd, Pr, Dy, Tb and Ho; T is one or more elements selected among Fe, Co and Ni; x is 0.13 or more but 0.30 or less; and y is 0.06 or more but 0.15 or less), the step (2) of heat-treating the cast alloy obtained in the above step (1) at 1,000 to 1,120°C, the step (3) of pulverizing the heat-treatable alloy obtained in the above step (2), and the step (4) of selectively collecting particles attracted by a permanent magnet or an electromagnet out of the powders obtained in the above step (3). <P>COPYRIGHT: (C)2004,JPO

Description

【０００１】
【発明の属する技術分野】
本発明は、焼結磁石用希土類合金粉末の製造方法に関するものであり、詳しくは、例えば家電品、ＯＡ機器、産業用機械などに使われる電磁型モータ用焼結磁石の原料となり得る焼結磁石用希土類合金粉末の製造方法に関するものである。
【０００２】
【従来の技術】
希土類−遷移金属−硼素系の焼結磁石は通常、溶解鋳造された合金を粉砕し、磁場中でプレス成形後、真空あるいは不活性ガス雰囲気中で焼結・熱処理することにより製造される。
図１は、希土類−遷移金属−硼素系鋳造合金の内部組織を示す模式図である。図１において、鋳造合金の内部組織は、その大部分が大きさ数十μｍのＲ_２Ｔ_１４Ｂ相１からなり、相境界部分に非磁性のＲ−ｒｉｃｈ相２やＢ−ｒｉｃｈ相３が分布している。これら非磁性相の量が多いと磁石の残留磁化を低下させる原因となるため、Ｒ−ｒｉｃｈ相やＢ−ｒｉｃｈ相はできる限り除去することが好ましい。その除去方法としては、例えば特開昭６４−４８４０６号公報に記載された技術が知られている。それによれば、鋳造合金に水素吸蔵放出処理を施して微粉末化した後、磁気選別を行うことにより、不要なＲ−ｒｉｃｈ相やＢ−ｒｉｃｈ相を除去している。
【０００３】
【発明が解決しようとする課題】
しかしながら、同公報の手法では磁気選別後に回収されたＲ_２Ｔ_１４Ｂ粒子の表面にはＲ−ｒｉｃｈ相が付着しており、その分のＲ−ｒｉｃｈ相は残存することになる。この点について同公報では一定量のＲ−ｒｉｃｈ相は保磁力を維持するために欠かせないとして容認している。ところが、Ｒ−ｒｉｃｈ相は酸素と結びつきやすいため、その後の磁石製造工程において粉末に不要な酸素が取り込まれてしまう。その結果、取り込まれた酸素が不純物として磁化に悪影響を及ぼす他、酸化による焼結性の低下などが起こり、最終的に残留磁化が低下してしまう問題点がある。同公報では、この問題点を見過ごしている。
【０００４】
したがって本発明の目的は、鋳造合金の粉末中に取り込まれる酸素量を低減し、不純物としての影響や焼結性の低下を回避し、残留磁化の低下を防ぐことができる焼結磁石用希土類合金粉末の製造方法の提供にある。
【０００５】
【課題を解決するための手段】
請求項１の発明は、
（１）Ｒ_ｘＴ_{１−ｘ−ｙ}Ｂ_ｙ合金（式中、ＲはＮｄ、Ｐｒ、Ｄｙ、ＴｂおよびＨｏより選ばれる１種以上、ＴはＦｅ、ＣｏおよびＮｉより選ばれる１種以上、０．１３≦ｘ≦０．３０、かつ０．０６≦ｙ≦０．１５である）を鋳造する工程と、
（２）前記（１）工程により得られた鋳造合金を１０００〜１１２０℃で熱処理する工程と、
（３）前記（２）工程により得られた熱処理合金を粉末化する工程と、
（４）前記（３）工程により得られた粉末から永久磁石あるいは電磁石により吸着する粒子を選択回収する工程と
を有することを特徴とする焼結磁石用希土類合金粉末の製造方法である。
請求項２の発明は、前記（３）工程には熱処理合金の水素吸蔵放出処理が含まれることを特徴とする請求項１に記載の焼結磁石用希土類合金粉末の製造方法である。
請求項３の発明は、前記（３）工程中、前記（４）工程中、前記（２）および（３）工程間、あるいは前記（４）工程後に、非晶質あるいは１００ｎｍ以下の結晶粒から構成される微結晶質のＲＭ−Ｔ−Ｂ合金（式中、ＲＭはＤｙ、Ｔｂのいずれか一方を必須とする希土類元素より選ばれる１種以上、ＴはＦｅ、ＣｏおよびＮｉより選ばれる１種以上である）を前記熱処理合金に対して５〜３０重量％の割合で混合することを特徴とする請求項１または２記載の焼結磁石用希土類合金粉末の製造方法である。
【０００６】
【作用】
図２は、本発明における（２）熱処理工程後の合金の内部組織を示す模式図である。図２において、前記組成合金を鋳造後１０００〜１１２０℃の熱処理を施すことにより合金内部のＲ_２Ｔ_１４Ｂ相１が２００μｍ以上の大きさに粗大化する。なお、符号２は、Ｒ−ｒｉｃｈ相であり、３はＢ−ｒｉｃｈ相である。
図３は、本発明における（３）粉末化工程によって派生する粒子を示す模式図である。粉末化工程を通じて熱処理合金の内部組織が破壊され、Ｒ_２Ｔ_１４Ｂ相１の内部よりＲ_２Ｔ_１４Ｂ粒子４Ａが、Ｒ_２Ｔ_１４Ｂ相１の外縁部よりＲ−ｒｉｃｈ相が一部付着したＲ_２Ｔ_１４Ｂ粒子４Ｂが、Ｒ−ｒｉｃｈ相よりＲ−ｒｉｃｈ粒子５が、Ｂ−ｒｉｃｈ相よりＢ−ｒｉｃｈ粒子６が派生する。ここで、Ｒ_２Ｔ_１４Ｂ相１は粗大化によりその比体積がかなり大きくなっているため、派生する粒子４Ａの数量は粒子４Ｂの１０倍以上となる。粉末化後はこれらの粒子が混在した状態であるが、永久磁石あるいは電磁石により吸着する粒子を取り出すと、強磁性であるＲ_２Ｔ_１４Ｂ粒子のみが選択的に回収される。こうして得られたＲ_２Ｔ_１４Ｂ粉末はその９割以上がＲ_２Ｔ_１４Ｂ粒子４Ａであり、その表面にはＲ−ｒｉｃｈ相がほとんど付着していないため、粉末に占めるＲ−ｒｉｃｈ相の体積比率が従来に比べて大幅に低減される。その結果、粉末に取り込まれる酸素量が激減する。なお、（３）粉末化工程では、１〜１５μｍの平均粒径まで粉末化するのが好ましい。
また、本発明において、熱処理合金を粉末化する工程の中に鋳造合金への水素吸蔵放出処理を加えることにより、合金の脆化が促進されて粉末化に要する時間が短縮化される。なお、水素吸蔵放出処理は、例えば常温常圧の水素気雰囲気中に数時間放置し、次いで１０^−２Ｔｏｒｒ以下の圧力下で３００〜６５０℃の温度で数時間加熱することにより行うことができる。
また、前記（３）工程中、前記（４）工程中、前記（２）および（３）工程間、あるいは前記（４）工程後に、非晶質あるいは１００ｎｍ以下の結晶粒から構成される微結晶質のＲＭ−Ｔ−Ｂ合金（式中、ＲＭはＤｙ、Ｔｂのいずれか一方を必須とする希土類元素より選ばれる１種以上、ＴはＦｅ、ＣｏおよびＮｉより選ばれる１種以上である）を焼結助剤として熱処理合金に対して５〜３０重量％の割合で混合する工程を加えることにより、焼結助剤がその後の磁石製造のための焼結・熱処理工程においてＲ−ｒｉｃｈ粒界相を形成するため、焼結性が上がると同時に、焼結助剤に含まれるＤｙあるいはＴｂがＲ_２Ｔ_１４Ｂ結晶表面に拡散することにより保磁力が大きく向上する。さらに焼結助剤は非晶質または微結晶質であり酸化されにくいため、混合しても粉末に不要な酸素が取り込まれない。
【０００７】
【発明の実施の形態】
実施の形態１．
高周波真空溶解炉にて種々の組成のＮｄ_ｘＦｅ_{１−ｘ−ｙ}Ｂ_ｙ合金（０．１０≦ｘ≦０．３５，０．０４≦ｙ≦０．１７）を溶解し、１３００〜１５００℃の温度に加熱後、水冷鋳型に鋳込んで厚み１０ｍｍの板状の鋳造合金を作製した。内部組織を光学顕微鏡で観察したところいずれの試料もＲ_２Ｔ_１４Ｂ相のサイズは３０μｍ程度であった。次に真空加熱炉を用いてこれらの合金を９００〜１１５０℃の種々の温度で８時間加熱した。その結果、１１２０℃を越える温度では大部分の鋳造合金がアルミナ製の台座と激しく反応して健全な熱処理が行われず、また、１０００℃を下回る温度ではＲ_２Ｔ_１４Ｂ相が５０μｍ程度までにしか成長せず、充分な粗大化が起こらないことがわかった。そのため、熱処理温度としては、１０００〜１１２０℃とすることが好ましい。同温度範囲で熱処理を行った場合の合金の内部組織について調べた結果を図４に示す。図中×印はα−Ｆｅ相、あるいはＮｄ_２Ｆｅ_１７相の存在が認められた組成であり、これらの相は磁石の磁気特性を大きく劣化させる要因となるため、原料にするには不適な組成である。また、△印は合金と炉内の台座との反応が激しく健全な熱処理が行えない組成、あるいはＲ_２Ｔ_１４Ｂ相の充分な粗大化が起こらない組成である。一方、○印は熱処理が健全に行われ２００μｍ以上のサイズのＲ_２Ｔ_１４Ｂ相が得られた組成で、この結果から合金組成は０．１３≦ｘ≦０．３０、０．０６≦ｙ≦０．１５の範囲が適当といえる。次に、Ｎｄ_０．１５Ｆｅ_０．７７Ｂ_０．０８、Ｎｄ_０．１２Ｐｒ_０．０３Ｆｅ_０．７７Ｂ_０．０８ _、Ｎｄ_０．１４Ｄｙ_０．０１Ｆｅ_０．７７Ｂ_０．０８、Ｎｄ_０．１４Ｔｂ_０．０１Ｆｅ_０．７７Ｂ_０．０８、Ｎｄ_０．１５Ｆｅ_０．７５Ｃｏ_０．０２Ｂ_０．０８、Ｎｄ_０．１５Ｆｅ_０．７６Ｎｉ_０．０１Ｂ_０．０８の６種類の合金試料を高周波真空溶解炉により鋳造した。その後真空加熱炉を用いて各合金試料を１１００℃の温度で８時間加熱した。内部組織を光学顕微鏡で観察したところいずれの試料もＲ_２Ｔ_１４Ｂ相が２００μｍ以上に粗大化していた。次にそれぞれの合金試料をジョークラッシャー、ディスクミル、ジェットミルにて順次粉砕し、得られた各粉末試料を永久磁石方式の磁力選別機にかけてＲ_２Ｔ_１４Ｂ粒子を選択回収した。このときＲ_２Ｔ_１４Ｂ粒子の平均粒径は４〜５μｍの範囲であった。また、粉体中のＲ−ｒｉｃｈ相、Ｂ−ｒｉｃｈ相の含有量はいずれも０．２体積％未満であり、Ｒ_２Ｆｅ_１４Ｂ相の含有量は９９．６体積％を超えることが確認された。一方で焼結助剤として液体急冷装置を用いてＮｄ_０．３８Ｆｅ_０．５６Ｂ_０．０６非晶質合金粉末を別途作製し、これを１２重量％の割合で各試料粉末に加えて均一に混合した後、磁場プレス機により圧力６００ｋｇｆ／ｃｍ^２、磁場１５０００Ｇの条件で横磁場成形を行った。次に、管状炉により真空中１０９０℃で２時間の焼結を行った後急冷し、さらに真空中で６００℃で２時間の熱処理を施し、外径１０ｍｍ×高さ７ｍｍの円柱磁石試料を得た。各試料を３００００Ｇのパルス磁場により着磁し、ＢＨカーブトレーサにより減磁曲線を測定した。また、不活性ガス融解−赤外線吸収法により含有酸素量を測定した。測定結果を表１に示す。同表には比較のため、鋳造後に１１００℃の加熱を行わず、焼結助剤も用いない従来の製法による磁石の特性を記載している。表１より本発明により得られた磁石の酸素濃度は１５００ｐｐｍ前後であり、比較例の磁石に比べて半分以下となっていることがわかる。これは本発明では合金粉末に含まれるＮｄ−ｒｉｃｈ相の量がわずかであり粉末の酸化量が低減されたことを反映している。一方、比較例では鋳造後に１１００℃の加熱を行っていないので、鋳造合金の結晶粒の粗大化が十分ではなく、このため、粉砕時には表面にＲ−ｒｉｃｈ相が付着したＲ_２Ｆｅ_１４Ｂ粒子が多数派生している。磁気選別後、回収された試料粉末に含まれるＲ−ｒｉｃｈ相の含有量を調べたところ、約３．０体積％と本発明よりはるかに多いことが分かった。この多量のＲ−ｒｉｃｈ相が製造工程中に多量の酸素を取り込んでしまい、その結果、比較例の含有酸素量は５０００ｐｐｍに近い高いレベルになってしまったといえる。次に、焼結・熱処理工程後の磁石の内部組織を調べたところ、本発明と比較例いずれもＲ−ｒｉｃｈ相の体積含有率は３．１％、Ｂ−ｒｉｃｈ相のそれは０．２％であった。本発明のＲ−ｒｉｃｈ相は主に焼結助剤が焼結工程中に相分離を起こして生成されるが、焼結助剤を非晶質もしくは微結晶質としているため、酸化されることがほとんどない。そのため、本発明ではＲ−ｒｉｃｈ相の酸化に起因する焼結性の低下がほとんど生じず、より緻密な内部組織を得ることができると同時に保持力の維持に必要な量のＲ−ｒｉｃｈ粒界相も良好に形成することができる。よって、含有酸素量の低減は、酸素が不純物として磁化へ悪影響を及ぼすことを回避するのみならず、焼結性の低下を防止する効果も有する。これらの相乗効果により本発明の磁石は表１に示すとおり、同組成を有する比較例の磁石と比べていずれも高い残留磁化が得られている。
【０００８】
【表１】

【０００９】
実施の形態２．
高周波真空溶解炉にてＮｄ_０．１５Ｆｅ_０．７７Ｂ_０．０８合金を溶解し１４５０℃の温度に加熱後、水冷鋳型に鋳込んで厚み１０ｍｍの板状の鋳造合金を作製した。内部組織を光学顕微鏡で観察したところＮｄ_２Ｆｅ_１４Ｂ相のサイズは３０μｍ程度であった。次に真空加熱炉を用いて同合金を１１００℃の温度で８時間加熱した。内部組織を観察したところＮｄ_２Ｆｅ_１４Ｂ相は２００μｍ以上に粗大化していた。次に同合金をジョークラッシャー、ディスクミル、ジェットミルにて順次粉砕して微粉末化し、次いで電磁石方式の磁力選別機にかけてＮｄ_２Ｆｅ_１４Ｂ粒子を選択回収して粉末試料を得た。同粉末試料の平均粒径は約５μｍであった。一方で焼結助剤として液体急冷装置を用いて約５０μｍ厚のＮｄ_０．２７Ｄｙ_０．１１Ｆｅ_０．５６Ｂ_０．０６合金薄帯を作製した。電子顕微鏡を用いて観察した結果、薄帯表面は非晶質であったが、内部約１５μｍ厚の部分は直径数十ｎｍ〜１００ｎｍの結晶粒が占めており、微結晶質となっていた。薄帯をディスクミル、ジェットミルにて順次粉砕して助剤粉末を得た。同粉末の平均粒径は約１μｍであった。粉末試料に助剤粉末を様々な割合で加えて均一に混合した後、磁場プレス機により圧力６００ｋｇｆ／ｃｍ^２、磁場１５０００Ｇの条件で横磁場成形を行った。次に、各成形体を管状炉により真空中１０９０℃で２時間の焼結を行った後急冷し、さらに真空中６００℃で２時間の熱処理を施し、外径１０ｍｍ×高さ７ｍｍの円柱磁石試料を得た。各試料を３００００Ｇのパルス磁場により着磁し、ＢＨカーブトレーサにより減磁曲線を測定した。測定結果を表２に示す。助剤粉末の粉末試料に対する重量比率が５％を下回ると保磁力が７ｋＯｅ未満となり、磁石の動作点によっては大きな減磁が起きる可能性が生じるため、信頼性に欠ける。また該重量比率が３０％を越えるとエネルギー積の低下が顕著になる。このため助剤粉末の粉末試料に対する重量比率は５〜３０％であることが好ましく、さらには１２〜２５％であることが好ましい。
【００１０】
【表２】

【００１１】
実施の形態３．
高周波真空溶解炉にてＮｄ_０．１５Ｆｅ_０．７８Ｂ_０．０７合金を溶解し１４５０℃の温度に加熱後、水冷鋳型に鋳込んで厚み１０ｍｍの板状の鋳造合金を作製した。内部組織を光学顕微鏡で観察したところＮｄ_２Ｆｅ_１４Ｂ相のサイズは３０μｍ程度であった。次に真空加熱炉を用いて同合金を１１００℃の温度で８時間加熱した。内部組織を観察したところＮｄ_２Ｆｅ_１４Ｂ相は２００μｍ以上に粗大化していた。次に同合金をジョークラッシャーで５ｍｍ角程度に粉砕して粗粉末を得た。一方で焼結助剤として液体急冷装置を用いてＮｄ_０．２７Ｔｂ_０．１１Ｆｅ_０．５６Ｂ_０．０６、Ｎｄ_０．２７Ｄｙ_０．１１Ｆｅ_０．５６Ｂ_０．０６、Ｐｒ_０．２７Ｄｙ_０．１１Ｆｅ_０．５６Ｂ_０．０６、Ｎｄ_０．３８Ｆｅ_０．５６Ｂ_０．０６、Ｐｒ_０．３８Ｆｅ_０．５６Ｂ_０．０６の５種類の合金薄帯を作製した。薄帯の厚みは約４０μｍであった。電子顕微鏡を用いて観察したところ薄帯内部は非晶質となっていた。粗粉末に対して各薄帯をそれぞれ１８重量％の重量比率で混合した後、各混合粉末を常温常圧の水素気雰囲気中に８時間放置し、継いで１０^−２Ｔｏｒｒの圧力下６００℃の温度で８時間加熱して水素放出を行った。その後ディスクミル、ジェットミルにて順次粉砕して微粉末を得た。ジェットミル粉砕に費やした時間は従来の１／４程度であったが、水素吸蔵放出処理による脆化効果により粉末は平均結晶粒径が４μｍ程度に微粒子化されていた。次に各微粉末を電磁石方式の磁力選別機にかけて粒子を回収した。このとき焼結助剤である非晶質合金薄帯から派生した微粒子も強磁性であるため電磁石に吸着し、Ｎｄ_２Ｆｅ_１４Ｂ粒子に混じって回収された。得られた各混合粉末を磁場プレス機により圧力６００ｋｇｆ／ｃｍ^２、磁場１５０００Ｇの条件で横磁場成形した。次に、各成形体を管状炉により真空中１０９０℃で２時間の焼結を行った後急冷し、さらに真空中６００℃で２時間の熱処理を施し、外径１０ｍｍ×高さ７ｍｍの円柱磁石試料を得た。各試料を３００００Ｇのパルス磁場により着磁し、ＢＨカーブトレーサにより減磁曲線を測定した。測定結果を表３に示す。ＤｙまたはＴｂを含む焼結助剤により作製した磁石はこれらの元素を含まない磁石に比べて大きな保磁力が得られた。
【００１２】
【表３】

【００１３】
【発明の効果】
請求項１の発明は、
（１）Ｒ_ｘＴ_{１−ｘ−ｙ}Ｂ_ｙ合金（式中、ＲはＮｄ、Ｐｒ、Ｄｙ、ＴｂおよびＨｏより選ばれる１種以上、ＴはＦｅ、ＣｏおよびＮｉより選ばれる１種以上、０．１３≦ｘ≦０．３０、かつ０．０６≦ｙ≦０．１５である）を鋳造する工程と、
（２）前記（１）工程により得られた鋳造合金を１０００〜１１２０℃で熱処理する工程と、
（３）前記（２）工程により得られた熱処理合金を粉末化する工程と、
（４）前記（３）工程により得られた粉末から永久磁石あるいは電磁石により吸着する粒子を選択回収する工程とを有することを特徴とする焼結磁石用希土類合金粉末の製造方法であるので、鋳造合金の粉末中に取り込まれる酸素量を低減し、不純物としての影響や焼結性の低下を回避し、残留磁化の低下を防ぐことができる。
【００１４】
請求項２の発明は、前記（３）工程には熱処理合金の水素吸蔵放出処理が含まれることを特徴とする請求項１に記載の焼結磁石用希土類合金粉末の製造方法であるので、合金の脆化が促進されて粉末化に要する時間が短縮化される。
【００１５】
請求項３の発明は、前記（３）工程中、前記（４）工程中、前記（２）および（３）工程間、あるいは前記（４）工程後に、非晶質あるいは１００ｎｍ以下の結晶粒から構成される微結晶質のＲＭ−Ｔ−Ｂ合金（式中、ＲＭはＤｙ、Ｔｂのいずれか一方を必須とする希土類元素より選ばれる１種以上、ＴはＦｅ、ＣｏおよびＮｉより選ばれる１種以上である）を前記熱処理合金に対して５〜３０重量％の割合で混合することを特徴とする請求項１または２記載の焼結磁石用希土類合金粉末の製造方法であるので、焼結性が上がると同時に保磁力も向上する。
【図面の簡単な説明】
【図１】希土類−遷移金属−硼素系鋳造合金の内部組織を示す模式図である。
【図２】本発明における熱処理後の希土類−遷移金属−硼素系鋳造合金の内部組織を示す模式図である。
【図３】本発明における熱処理後の希土類−遷移金属−硼素系鋳造合金の粉砕よって派生する粒子を示す模式図である。
【図４】実施の形態１において健全な特性を有する粉末が得られる組成領域を示した図である。
【符号の説明】
１　Ｒ_２Ｔ_１４Ｂ相、２　Ｒ−ｒｉｃｈ相、３　Ｂ−ｒｉｃｈ相、４Ａ　Ｒ_２Ｔ_１４Ｂ相内部より派生したＲ_２Ｔ_１４Ｂ粒子、４Ｂ　Ｒ_２Ｔ_１４Ｂ相外縁部より派生したＲ_２Ｔ_１４Ｂ粒子、５　Ｒ−ｒｉｃｈ相より派生したＲ−ｒｉｃｈ粒子、６　Ｂ−ｒｉｃｈ相より派生したＢ−ｒｉｃｈ粒子。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a method for producing a rare earth alloy powder for a sintered magnet, and more specifically, a sintered magnet that can be used as a raw material for a sintered magnet for an electromagnetic motor used for home appliances, OA equipment, industrial machinery, and the like. The present invention relates to a method for producing a rare earth alloy powder for use.
[0002]
[Prior art]
Rare earth-transition metal-boron based sintered magnets are usually produced by pulverizing a melt-cast alloy, press-molding in a magnetic field, and then sintering and heat-treating in a vacuum or inert gas atmosphere.
FIG. 1 is a schematic diagram showing the internal structure of a rare earth-transition metal-boron cast alloy. In FIG. 1, most of the internal structure of the cast alloy is composed of an R ₂ T ₁₄ B phase 1 having a size of several tens of μm, and a nonmagnetic R-rich phase 2 and a B-rich phase 3 are formed at a phase boundary. Are distributed. If the amount of these non-magnetic phases is large, the residual magnetization of the magnet may be reduced. Therefore, it is preferable to remove the R-rich phase and the B-rich phase as much as possible. As a removing method, for example, a technique described in JP-A-64-48406 is known. According to this method, unnecessary R-rich phase and B-rich phase are removed by subjecting the cast alloy to a hydrogen storage / release treatment to make it into fine powder and then performing magnetic separation.
[0003]
[Problems to be solved by the invention]
However, according to the method disclosed in the publication, the R-rich phase is attached to the surface of the R ₂ T ₁₄ B particles collected after the magnetic separation, and the R-rich phase remains there. In this regard, the publication accepts that a certain amount of the R-rich phase is indispensable for maintaining coercive force. However, since the R-rich phase is easily linked to oxygen, unnecessary oxygen is taken into the powder in the subsequent magnet manufacturing process. As a result, there is a problem that the taken-in oxygen adversely affects the magnetization as an impurity, and the sinterability is reduced due to oxidation, and the residual magnetization is eventually reduced. The publication overlooks this problem.
[0004]
Therefore, an object of the present invention is to reduce the amount of oxygen taken into the powder of a cast alloy, to avoid the influence of impurities and a decrease in sinterability, and to prevent a decrease in remanent magnetization. An object of the present invention is to provide a method for producing a powder.
[0005]
[Means for Solving the Problems]
The invention of claim 1 is
(1) R _x T _1- _{xy By} alloy (where R is at least one selected from Nd, Pr, Dy, Tb and Ho, T is at least one selected from Fe, Co and Ni, 0.13 ≦ x ≦ 0.30, and 0.06 ≦ y ≦ 0.15).
(2) a step of heat-treating the cast alloy obtained in the step (1) at 1000 to 1120 ° C;
(3) a step of powdering the heat-treated alloy obtained in the step (2);
(4) a step of selectively recovering particles adsorbed by a permanent magnet or an electromagnet from the powder obtained in the above step (3).
The invention according to claim 2 is the method for producing a rare earth alloy powder for a sintered magnet according to claim 1, wherein the step (3) includes a hydrogen storage / release treatment of the heat-treated alloy.
The invention according to claim 3 is that, during the step (3), during the step (4), between the steps (2) and (3), or after the step (4), amorphous or crystal grains of 100 nm or less are formed. A microcrystalline RM-TB alloy (where RM is at least one element selected from the group consisting of Dy and Tb, and T is 1 element selected from Fe, Co and Ni). The rare-earth alloy powder for a sintered magnet according to claim 1 or 2, wherein 5% by weight or more is mixed with the heat-treated alloy at a ratio of 5 to 30% by weight.
[0006]
[Action]
FIG. 2 is a schematic view showing the internal structure of the alloy after the heat treatment step (2) in the present invention. In FIG. 2, the R ₂ T ₁₄ B phase 1 inside the alloy is coarsened to a size of 200 μm or more by performing a heat treatment at 1000 to 1120 ° C. after casting the composition alloy. Note that reference numeral 2 denotes an R-rich phase, and reference numeral 3 denotes a B-rich phase.
FIG. 3 is a schematic diagram showing particles derived by the (3) powdering step in the present invention. Internal structure of the heat-treated alloy through powdered process is _disrupted, R _{2 T 14} inside than _R 2 _{T 14} B grains 4A of B phase _1, R _{2 T 14} R-rich phase is a part of the outer edge portion of the B-phase 1 The attached R ₂ T ₁₄ B particles 4B, R-rich particles 5 are derived from the R-rich phase, and B-rich particles 6 are derived from the B-rich phase. Here, since the specific volume of the R ₂ T ₁₄ B phase 1 is considerably large due to the coarsening, the number of the derived particles 4A is ten times or more the number of the particles 4B. After powdering, these particles are in a mixed state. However, when particles adsorbed by a permanent magnet or an electromagnet are taken out, only ferromagnetic R ₂ T ₁₄ B particles are selectively recovered. 90% or more of the R ₂ T ₁₄ B powder thus obtained is R ₂ T ₁₄ B particles 4A, and the R-rich phase hardly adheres to the surface thereof. The volume ratio is greatly reduced as compared with the conventional case. As a result, the amount of oxygen taken into the powder is drastically reduced. In the powdering step (3), powdering is preferably performed to an average particle diameter of 1 to 15 μm.
Further, in the present invention, by adding a hydrogen absorbing and releasing treatment to the cast alloy during the step of powdering the heat-treated alloy, embrittlement of the alloy is promoted and the time required for powdering is shortened. The hydrogen storage and release treatment can be performed, for example, by leaving the substrate in an atmosphere of hydrogen at normal temperature and normal pressure for several hours, and then heating at a temperature of 300 to 650 ° C. under a pressure of 10 ⁻² Torr or less for several hours. .
Further, in the step (3), in the step (4), between the steps (2) and (3), or after the step (4), amorphous or microcrystals composed of crystal grains of 100 nm or less. RM-TB alloy (wherein, RM is at least one selected from rare earth elements in which one of Dy and Tb is essential, and T is at least one selected from Fe, Co and Ni) Is added as a sintering aid to the heat-treated alloy at a ratio of 5 to 30% by weight, so that the sintering aid can be used in the subsequent sintering and heat-treating step for magnet production in the R-rich grain boundary. Since a phase is formed, sinterability is improved, and at the same time, Dy or Tb contained in the sintering aid is diffused to the R ₂ T ₁₄ B crystal surface, thereby greatly improving coercive force. Further, since the sintering aid is amorphous or microcrystalline and is hardly oxidized, unnecessary oxygen is not taken into the powder even when mixed.
[0007]
BEST MODE FOR CARRYING OUT THE INVENTION
Embodiment 1 FIG.
In a high frequency vacuum melting furnace, Nd _x Fe _1- _{xy By} alloys of various compositions (0.10 ≦ x ≦ 0.35, 0.04 ≦ y ≦ 0.17) are melted, and 1300-1500 ° C. After heating to a temperature of, a cast alloy having a thickness of 10 mm was cast into a water-cooled mold. When the internal structure was observed with an optical microscope, the size of the R ₂ T ₁₄ B phase was about 30 μm in all samples. Next, these alloys were heated at various temperatures of 900 to 1150 ° C. for 8 hours using a vacuum heating furnace. As a result, at temperatures above 1120 ° C. Most of the cast alloy is not performed healthy heat treatment reacts violently with alumina pedestal, also before 50μm approximately the R _{2 T} 14 _B phase at temperatures below 1000 ° C. However, it was found that sufficient growth did not occur. Therefore, the heat treatment temperature is preferably set to 1000 to 1120 ° C. FIG. 4 shows the results of examination of the internal structure of the alloy when heat treatment was performed in the same temperature range. The crosses in the figure indicate the composition in which the presence of the α-Fe phase or the Nd ₂ Fe ₁₇ phase was recognized, and these phases are factors that greatly deteriorate the magnetic properties of the magnet, and are unsuitable as raw materials. The composition. The symbol “△” represents a composition in which the reaction between the alloy and the pedestal in the furnace is so severe that a sound heat treatment cannot be performed, or a composition in which the R ₂ T ₁₄ B phase does not sufficiently grow. On the other hand, the mark “○” indicates a composition in which heat treatment was performed properly and an R ₂ T ₁₄ B phase having a size of 200 μm or more was obtained. From these results, the alloy compositions were 0.13 ≦ x ≦ 0.30 and 0.06 ≦ y. It can be said that the range of ≦ 0.15 is appropriate. Next, Nd _0.15 Fe _0.77 B _0.08 , Nd _0.12 Pr _0.03 Fe _0.77 B _0.08 _, Nd _0.14 Dy _0.01 Fe _0.77 B _0.08 , Nd _0.14 Tb _0.01 Fe _0.77 B _0.08 , Nd _0.15 Fe _0.75 Co _0.02 B _0.08 , Nd _0.15 Fe _0.76 Ni _0.01 B _0.08 Were cast in a high-frequency vacuum melting furnace. Thereafter, each alloy sample was heated at 1100 ° C. for 8 hours using a vacuum heating furnace. When the internal structure was observed with an optical microscope, the R ₂ T ₁₄ B phase of each sample was coarsened to 200 μm or more. Next, each alloy sample was sequentially pulverized by a jaw crusher, a disc mill, and a jet mill, and each of the obtained powder samples was passed through a permanent magnet type magnetic separator to selectively recover R ₂ T ₁₄ B particles. At this time, the average particle size of the R ₂ T ₁₄ B particles was in the range of 4 to 5 μm. The content of the R-rich phase and the content of the B-rich phase in the powder were both less than 0.2% by volume, and the content of the R ₂ Fe ₁₄ B phase exceeded 99.6% by volume. Was done. On the other hand, an Nd _0.38 Fe _0.56 B _0.06 amorphous alloy powder was separately prepared using a liquid quenching device as a sintering aid, and this was added to each sample powder at a rate of 12% by weight to be uniform. After that, a horizontal magnetic field was formed by a magnetic field press under the conditions of a pressure of 600 kgf / cm ² and a magnetic field of 15000G. Next, after sintering at 1090 ° C. for 2 hours in a vacuum using a tube furnace, rapid cooling is performed, and heat treatment is further performed at 600 ° C. for 2 hours in a vacuum to obtain a cylindrical magnet sample having an outer diameter of 10 mm and a height of 7 mm. Was. Each sample was magnetized by a pulse magnetic field of 30,000 G, and a demagnetization curve was measured by a BH curve tracer. The oxygen content was measured by an inert gas melting-infrared absorption method. Table 1 shows the measurement results. In the same table, for comparison, the characteristics of a magnet manufactured by a conventional method without heating at 1100 ° C. after casting and using no sintering aid are described. From Table 1, it can be seen that the oxygen concentration of the magnet obtained according to the present invention is around 1500 ppm, which is less than half that of the magnet of the comparative example. This reflects that in the present invention, the amount of the Nd-rich phase contained in the alloy powder was small, and the oxidation amount of the powder was reduced. On the other hand, since no then heated for 1100 ° C. after casting in the comparative example, not enough grain coarsening of the cast alloy, Therefore, R-rich phase on the surface at the time of pulverization is attached R ₂ Fe 14 _B particles Has been derived in large numbers. After the magnetic separation, the content of the R-rich phase contained in the recovered sample powder was examined. As a result, it was found that the content was about 3.0% by volume, which was much higher than that of the present invention. It can be said that this large amount of the R-rich phase took in a large amount of oxygen during the manufacturing process, and as a result, the oxygen content of the comparative example reached a high level close to 5000 ppm. Next, when the internal structure of the magnet after the sintering / heat treatment step was examined, the volume content of the R-rich phase was 3.1% and that of the B-rich phase was 0.2% in both the present invention and the comparative example. Met. The R-rich phase of the present invention is mainly produced by causing the sintering aid to undergo phase separation during the sintering step, but is oxidized because the sintering aid is amorphous or microcrystalline. There is almost no. Therefore, in the present invention, a decrease in sinterability due to oxidation of the R-rich phase hardly occurs, and a more dense internal structure can be obtained, and at the same time, an amount of R-rich grain boundary necessary for maintaining a holding force is obtained. The phase can also be formed well. Therefore, the reduction of the oxygen content not only prevents the oxygen from adversely affecting the magnetization as an impurity, but also has the effect of preventing a decrease in sinterability. As shown in Table 1, due to the synergistic effects, the magnets of the present invention have higher remanence as compared with the magnet of the comparative example having the same composition.
[0008]
[Table 1]

[0009]
Embodiment 2 FIG.
A Nd _0.15 Fe _0.77 B _0.08 alloy was melted in a high-frequency vacuum melting furnace, heated to a temperature of 1450 ° C., and then cast into a water-cooled mold to produce a plate-shaped cast alloy having a thickness of 10 mm. When the internal structure was observed with an optical microscope, the size of the Nd ₂ Fe ₁₄ B phase was about 30 μm. Next, the alloy was heated at 1100 ° C. for 8 hours using a vacuum heating furnace. When the internal structure was observed, the Nd ₂ Fe ₁₄ B phase was coarsened to 200 μm or more. Next, the same alloy was sequentially pulverized with a jaw crusher, a disc mill, and a jet mill to finely powder, and then subjected to an electromagnet type magnetic separator to selectively collect Nd ₂ Fe ₁₄ B particles to obtain a powder sample. The average particle size of the powder sample was about 5 μm. On the other hand, an Nd _0.27 Dy _0.11 Fe _0.56 B _0.06 alloy ribbon having a thickness of about 50 μm was produced using a liquid quenching device as a sintering aid. As a result of observation using an electron microscope, the surface of the ribbon was amorphous, but a portion having a thickness of about 15 μm was occupied by crystal grains having a diameter of several tens nm to 100 nm, and was microcrystalline. The ribbon was pulverized sequentially with a disk mill and a jet mill to obtain an auxiliary powder. The average particle size of the powder was about 1 μm. Auxiliary powder was added to the powder sample at various ratios and mixed uniformly, and then subjected to transverse magnetic field molding using a magnetic field press under the conditions of a pressure of 600 kgf / cm ² and a magnetic field of 15000G. Next, each compact was sintered in a tube furnace at 1090 ° C. for 2 hours in a vacuum, then quenched, and further subjected to a heat treatment at 600 ° C. for 2 hours in a vacuum to obtain a cylindrical magnet having an outer diameter of 10 mm and a height of 7 mm. A sample was obtained. Each sample was magnetized by a pulse magnetic field of 30,000 G, and a demagnetization curve was measured by a BH curve tracer. Table 2 shows the measurement results. If the weight ratio of the auxiliary powder to the powder sample is less than 5%, the coercive force becomes less than 7 kOe, and there is a possibility that a large demagnetization may occur depending on the operating point of the magnet, so that the reliability is lacking. When the weight ratio exceeds 30%, the energy product is significantly reduced. Therefore, the weight ratio of the auxiliary powder to the powder sample is preferably 5 to 30%, and more preferably 12 to 25%.
[0010]
[Table 2]

[0011]
Embodiment 3 FIG.
A Nd _0.15 Fe _0.78 B _0.07 alloy was melted in a high-frequency vacuum melting furnace, heated to a temperature of 1450 ° C., and then cast into a water-cooled mold to produce a plate-shaped cast alloy having a thickness of 10 mm. When the internal structure was observed with an optical microscope, the size of the Nd ₂ Fe ₁₄ B phase was about 30 μm. Next, the alloy was heated at 1100 ° C. for 8 hours using a vacuum heating furnace. When the internal structure was observed, the Nd ₂ Fe ₁₄ B phase was coarsened to 200 μm or more. Next, the alloy was ground to about 5 mm square with a jaw crusher to obtain a coarse powder. On the other hand, Nd _0.27 Tb _0.11 Fe _0.56 B _0.06 , Nd _0.27 Dy _0.11 Fe _0.56 B _0.06 , Pr _0. Five kinds of alloy ribbons of ₂₇ Dy _0.11 Fe _0.56 B _0.06 , Nd _0.38 Fe _0.56 B _0.06 , and Pr _0.38 Fe _0.56 B _0.06 were produced. The thickness of the ribbon was about 40 μm. Observation using an electron microscope revealed that the inside of the ribbon was amorphous. After mixing each of the ribbons with the coarse powder at a weight ratio of 18% by weight, each of the mixed powders was allowed to stand in a hydrogen atmosphere at normal temperature and normal pressure for 8 hours, and then joined at 600 ° C. under a pressure of 10 ⁻² Torr. For 8 hours to release hydrogen. Thereafter, the resultant was pulverized sequentially with a disk mill and a jet mill to obtain fine powder. The time spent in jet mill pulverization was about one-fourth of the conventional one, but the powder was finely divided into particles having an average crystal grain size of about 4 μm due to the embrittlement effect of the hydrogen absorbing and releasing treatment. Next, each fine powder was passed through an electromagnet type magnetic force separator to collect particles. At this time, the fine particles derived from the amorphous alloy ribbon, which is a sintering aid, were also ferromagnetic and were adsorbed by the electromagnet, and were collected in the Nd ₂ Fe ₁₄ B particles. Each of the obtained mixed powders was subjected to a transverse magnetic field molding under a condition of a pressure of 600 kgf / cm ² and a magnetic field of 15000 G by a magnetic field press. Next, each compact was sintered in a tube furnace at 1090 ° C. for 2 hours in a vacuum, then quenched, and further subjected to a heat treatment at 600 ° C. for 2 hours in a vacuum to obtain a cylindrical magnet having an outer diameter of 10 mm and a height of 7 mm. A sample was obtained. Each sample was magnetized by a pulse magnetic field of 30,000 G, and a demagnetization curve was measured by a BH curve tracer. Table 3 shows the measurement results. Magnets made with sintering aids containing Dy or Tb provided greater coercive force than magnets without these elements.
[0012]
[Table 3]

[0013]
【The invention's effect】
The invention of claim 1 is
(1) R _x T _1- _{xy By} alloy (where R is at least one selected from Nd, Pr, Dy, Tb and Ho, T is at least one selected from Fe, Co and Ni, 0.13 ≦ x ≦ 0.30, and 0.06 ≦ y ≦ 0.15).
(2) a step of heat-treating the cast alloy obtained in the step (1) at 1000 to 1120 ° C;
(3) a step of powdering the heat-treated alloy obtained in the step (2);
And (4) a step of selectively recovering particles adsorbed by a permanent magnet or an electromagnet from the powder obtained in the step (3). It is possible to reduce the amount of oxygen taken into the powder of the alloy, to avoid the influence of impurities and a decrease in sinterability, and to prevent a decrease in residual magnetization.
[0014]
The invention according to claim 2 is the method for producing a rare earth alloy powder for a sintered magnet according to claim 1, wherein the step (3) includes a hydrogen storage / release treatment of the heat-treated alloy. And the time required for powdering is shortened.
[0015]
The invention according to claim 3 is that, during the step (3), during the step (4), between the steps (2) and (3), or after the step (4), amorphous or crystal grains of 100 nm or less are formed. A microcrystalline RM-TB alloy (where RM is at least one element selected from the group consisting of Dy and Tb, and T is 1 element selected from Fe, Co and Ni). A rare earth alloy powder for a sintered magnet according to

claim

1 or 2, wherein the heat-treated alloy is mixed at a ratio of 5 to 30% by weight with respect to the heat-treated alloy. At the same time, the coercive force also increases.
[Brief description of the drawings]
FIG. 1 is a schematic diagram showing the internal structure of a rare earth-transition metal-boron cast alloy.
FIG. 2 is a schematic diagram showing an internal structure of a rare earth-transition metal-boron cast alloy after heat treatment in the present invention.
FIG. 3 is a schematic view showing particles derived by grinding of a rare earth-transition metal-boron cast alloy after heat treatment in the present invention.
FIG. 4 is a diagram showing a composition region in which powder having sound characteristics is obtained in the first embodiment.
[Explanation of symbols]
1 R ₂ T ₁₄ B phase, 2 R-rich phase, 3 B-rich phase, _4A R _{2 T 14} B phase inside than derived _R 2 _{T 14} B grains, derived from _4B R _{2 T 14} B phase outer edge R ₂ T ₁₄ B particles, 5 R-rich particles derived from the R-rich phase, 6 B-rich particles derived from the B-rich phase.

Claims

(1) R _x T _1- _{xy By} alloy (where R is at least one selected from Nd, Pr, Dy, Tb and Ho, T is at least one selected from Fe, Co and Ni, 0.13 ≦ x ≦ 0.30, and 0.06 ≦ y ≦ 0.15).
(2) a step of heat-treating the cast alloy obtained in the step (1) at 1000 to 1120 ° C;
(3) a step of powdering the heat-treated alloy obtained in the step (2);
And (4) a step of selectively recovering particles adsorbed by a permanent magnet or an electromagnet from the powder obtained in the step (3).

The method for producing a rare earth alloy powder for a sintered magnet according to claim 1, wherein the step (3) includes a process of absorbing and releasing hydrogen from the heat-treated alloy.

During the step (3), during the step (4), between the steps (2) and (3), or after the step (4), an amorphous or microcrystalline material composed of crystal grains of 100 nm or less. An RM-TB alloy (wherein, RM is at least one member selected from rare earth elements that require one of Dy and Tb, and T is at least one member selected from Fe, Co, and Ni) The method for producing a rare earth alloy powder for a sintered magnet according to claim 1 or 2, wherein the mixture is mixed at a ratio of 5 to 30% by weight with respect to the heat-treated alloy.