JP3586577B2

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Info

Publication number: JP3586577B2
Application number: JP02392899A
Authority: JP
Inventors: 公穂内田; 昌弘高橋; 文丈谷口
Original assignee: Neomax Co Ltd
Current assignee: Hitachi Metals Ltd
Priority date: 1999-02-01
Filing date: 1999-02-01
Publication date: 2004-11-10
Anticipated expiration: 2016-03-19
Also published as: JPH11273922A

Description

【０００１】
【発明の属する技術分野】
本発明は、R-Fe-B系の焼結型永久磁石の性能改善に関するものである。
【０００２】
【従来の技術】
焼結型希土類永久磁石の中でＲ−Ｆｅ−Ｂ系（ＲはＹを含む希土類元素のうちの１種又は２種以上）焼結型永久磁石は高性能磁石として注目され、広い分野で使用されている。
このＲ−Ｆｅ−Ｂ系焼結型永久磁石は、基本的にはＲ_２Ｆｅ_１４Ｂ相（主相）、ＲＦｅ_７Ｂ_６相（Ｂｒｉｃｈ相）、Ｒ_８５Ｆｅ_１５相（Ｒｒｉｃｈ相）の３相からなる構造を有している。組成的に希土類元素に豊んだＲｒｉｃｈ相の存在と、このような３相構造に由来して、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石はＳｍ−Ｃｏ系焼結型永久磁石に比べて耐蝕性が劣り、この永久磁石の開発当初から現在に至るまで欠点の１つとなっている。
Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の腐蝕のメカニズムについての定説は無いが、Ｒｒｉｃｈ相を起点とした腐蝕の形態が一般的であることから、Ｒｒｉｃｈ相を陽極とした陽極腐蝕との見方もある。確かに、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の希土類元素の量を減少することによって、その焼結体内部のＲｒｉｃｈ相の量は減少し、かつ相の形態は微細化し、これに対応して永久磁石の耐蝕性は向上する。従って、希土類元素の量を減少することは、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の耐蝕性改善の一つの方法である。
【０００３】
Ｒ−Ｆｅ−Ｂ系を含む焼結型の希土類永久磁石は、原料金属を溶解し鋳型に注湯して得られたインゴットを粉砕，成形，焼結，熱処理，加工するという粉末冶金的な工程によって製造されるのが一般的である。しかし、インゴットを粉砕して得られる合金粉末は、希土類元素を多量に含むため化学的に非常に活性であり、大気中において酸化して含有酸素量が増加する。これによって、得られた焼結体の希土類元素の一部が酸化物を形成し、磁気的に有効な希土類元素が減少する。このため、実用的な磁気特性の水準、例えばｉＨｃ≧１３ｋＯｅを実現するためには、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の希土類元素の量を増やす必要があり、重量百分比率で３１％を越える希土類元素の添加量が実用材料では採用されている。
このため、これまでのＲ−Ｆｅ−Ｂ系焼結型永久磁石の耐蝕性は十分ではなかった。
【０００４】
【発明が解決しようとする課題】
本発明は、以上述べたR-Fe-B系焼結型永久磁石の耐蝕性を大幅に改善しようとするものである。
【０００５】
【課題を解決するための手段】
本発明者らは、R-Fe-B系焼結型永久磁石の耐蝕性を改善するため種々検討した結果、特定範囲量の希土類量と特定量以下の酸素量と炭素量のR-Fe-B系焼結型永久磁石において、その含有窒素量を特定範囲量とすることによって、耐蝕性が改善されるとともに実用的な高い磁気特性が得られることを見い出した。そして、R-Fe-B系焼結型永久磁石の耐蝕性をさらに改善すべく研究を継続した結果、その磁石主相結晶粒径を特定値以下とすることによって、耐蝕性がさらに向上することを見い出して本発明に至ったものである。
【０００６】
以下、本発明を具体的に説明する。
本発明の焼結型永久磁石は、重量百分率でR（RはYを含む希土類元素のうちの１種又は２種以上）27.0〜31.0％,B 0.5〜2.0％,N 0.02〜0.15％,O 0.25％以下,C 0.15％以下,残部Feの組成を有する焼結型永久磁石であって、実質的に前記焼結型永久磁石の主相およびRrich相に相当する２相組織を有する前記焼結型永久磁石用のR-Fe-B系急冷合金を用いて製造され、耐蝕性の向上していることを特徴とする。
また本発明の焼結型永久磁石は、重量百分率でR（RはYを含む希土類元素のうちの１種又は２種以上）27.0〜31.0％,B 0.5〜2.0％,N 0.02〜0.15％,O 0.25％以下,C 0.15％以下,残部Feの組成を有し、磁石主相結晶粒の総面積に対し、結晶粒径が10μｍ以下の主相結晶粒の面積の和が80％以上,結晶粒径が13μｍ以上の主相結晶粒の面積の和が10％以下である焼結型永久磁石であって、実質的に前記焼結型永久磁石の主相およびRrich相に相当する２相組織を有する前記焼結型永久磁石用のR-Fe-B系急冷合金を用いて製造され、耐蝕性の向上していることを特徴とする。
実質的に前記焼結型永久磁石の主相およびRrich相に相当する２相組織を有する前記焼結型永久磁石用のR-Fe-B系急冷合金は、いわゆるストリップキャスト法により製造される板厚0.1〜0.5mmの急冷薄帯状合金であり、本発明の焼結型永久磁石の主相に相当する相、Rrich相に相当する相以外に少量のBrich相を含むが実質的に前記の主相に相当する相およびRrich相に相当する相の２相からなる。
前記焼結型永久磁石において、Feの一部をNb 0.1〜2.0％,Al 0.02〜2.0％,Co 0.3〜5.0％,Ga 0.01〜0.5％,Cu 0.01〜1.0％のうちの１種又は２種以上で置換することができる。
【０００７】
本発明者らは、上記組成を有するＲ−Ｆｅ−Ｂ系焼結型永久磁石の耐蝕性に結晶粒径依存性があり、磁石主相結晶粒径を特定値以下にすることによって、特に優れた耐蝕性が発現されることを見い出した。磁石結晶粒径の定義と測定には種々の方法があり得、一義的ではないが、本発明者らは磁石主相の総面積に対する粒径が一定寸法以下の主相結晶粒の面積の和の割合と、同じく磁石主相の総面積に対する粒径が一定寸法以上の主相結晶粒の面積の和の割合によって、磁石結晶粒径の状態を示す尺度とした。以下この尺度を用いて本発明の効果を説明することとする。また、この割合を算出するに当たっての計測は、対象とするＲ−Ｆｅ−Ｂ系焼結型永久磁石の結晶組織を、ＯＬＹＭＰＵＳ社製顕微鏡（商品名ＶＡＮＯＸ）で観察し、この画像をＮＩＲＥＣＯ社製画像処理装置（商品名ＬＵＺＥＸ２）に直接投入して行った。
【０００８】
図１は、重量百分率でＮｄ２７．５％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．１％，Ａｌ０．１％，Ｃｏ２．０％，Ｇａ０．０８％，Ｎ０．０６％，Ｏ０．１６％，Ｃ０．０６％，Ｎ０．０４０％，残部Ｆｅの組成を有し、磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が９６％，結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１％の焼結型永久磁石の光学顕微鏡（１０００倍で観察）による観察結果である。また図２は、同じ組成を有し、磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が６４％，結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１７％の焼結型永久磁石の光学顕微鏡（１０００倍で観察）による観察結果である。
【０００９】
これらの焼結型永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの寸法に加工後、その表面に膜厚２０μｍのＮｉメッキを施した。次いで、試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べた。結晶粒径が小さな図１の永久磁石から作製した試料では、２５００Ｈｒ経過でもＮｉメッキに異常が見られなかった。５０００Ｈｒの経過では、Ｎｉメッキのわずかなハク離が観察された。一方、比較的大きな結晶粒が存在する図２の永久磁石から作製した試料では、１０００Ｈｒ経過でもＮｉメッキに異常が見られなかった。上記評価方法は加速試験であるため、この結果から、図２の永久磁石の耐蝕性は実用上全く問題ないと言える。しかし、２０００Ｈｒの経過においてはこの試料にＮｉメッキの大きなハク離が観察され、このことから図１の永久磁石と図２の永久磁石の間には厳密には耐蝕性に差があることが判った。即ち、磁石主相結晶粒径が小さいほど耐蝕性は良好である。図３は、図１の永久磁石から作製した試料の、上記評価試験を５０００Ｈｒおこなった後の断面のＳＥＭによる観察結果である。Ｎｉメッキと下地である永久磁石焼結体との間に部分的なハク離はあるものの、両者の密着性は比較的良好である。また、５０００Ｈｒの加速試験によっても、下地である永久磁石焼結体はほとんど損傷を受けていないことがわかる。
図４は、図２の試料から作製した試料の、上記評価試験を２０００Ｈｒおこなった後の断面の、ＳＥＭによる観察結果である。加速試験によって、下地である永久磁石焼結体の結晶粒界自体が破壊され、これによってＮｉメッキの大きなハク離が生じていることがわかる。
【００１０】
以上の結果から、永久磁石焼結体の主相結晶粒径の違いによって、耐蝕性の加速試験に対する結晶粒界の破壊のされ方に差異があることがわかった。この原因を図４から推定すると、比較的大きな主相結晶粒が存在する図２のような永久磁石焼結体においては、相対的に主相結晶粒の間の空隙部、具体的には粒界３重点がその主たる部分であり、ここにはきわめて酸化されやすいＮｄｒｉｃｈ相が存在しているが、このＮｄｒｉｃｈ相で充填されている空隙部の体積が大きくなる。腐蝕破壊をもたらす因子、例えば本加速試験では水分であるが、この様な因子の浸透性が良く、結晶粒界の破壊が連鎖反応的に起こりやすい状態にあるものと考えられる。以上は、本発明に係るＲ−Ｆｅ−Ｂ系焼結型永久磁石の耐蝕性に主相結晶粒径依存性があることを、本発明者らの研究結果の一例を示すことによって説明したものである。
【００１１】
本発明者らは、本発明で採用したＲ−Ｆｅ−Ｂ系焼結型永久磁石の組成、主相結晶粒径分布と、磁気特性、耐蝕性との相関について上記の様な評価を継続し、図５に示すような結果を得た。
図５は、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和の割合と、同じく磁石主相結晶の総面積に対する結晶粒径が１３μｍ以上の主相の結晶粒の面積の和の割合と、耐蝕性の加速試験での、Ｎｉメッキのハク離開始が生じるまでの経過時間との関係を示したものである。○印は重量百分比率でＮｄ２５．５％，Ｐｒ２．５％，Ｄｙ２．０％，Ｂ１．０％，Ｎｂ０．４％，Ａｌ１．０％，Ｃｏ３．０％，Ｃｕ０．１％，Ｏ０．１９％，Ｃ０．０８％，Ｎ０．０４０％，残部Ｆｅの組成を有する焼結体、□印は重量百分比率でＮｄ２８．０％，Ｄｙ１．０％，Ｂ１．０５％，Ａｌ０．０５％，Ｃｏ２．０％，Ｇａ０．０９％，Ｏ０．２０％，Ｃ０．０７％，Ｎ０．０８０％，残部Ｆｅの組成を有する焼結体、△印は重量百分比率でＮｄ２４．５％，Ｐｒ１．５％，Ｄｙ４．５％，Ｂ１．１％，Ｎｂ１．０％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．１％，Ｃｕ０．０８％，Ｏ０．１８％，Ｃ０．０６％，Ｎ０．０５０％，残部Ｆｅの組成を有する焼結体を示す。この場合の加速試験では、磁石を８ｍｍ×８ｍｍ×２ｍｍの寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施し、次いで試料を２気圧，１２０℃，湿度１００％の条件に放置した。図５から、磁石主相の結晶の総面積に対し、結晶粒径が１０μｍ以下の主相結晶粒の面積の和が８０％以上で、かつ結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１０％以下である場合において、前記特定組成を有するＲ−Ｆｅ−Ｂ系焼結型永久磁石の耐蝕性が特に優れたものになることがわかる。従って、磁石主相結晶粒の大きさは、上記に規定される。
【００１２】
前記特定組成を有するＲ−Ｆｅ−Ｂ系焼結型永久磁石の主相の結晶粒径を上記の規定範囲のものに制御する方法は必ずしも一義的ではなく、種々の方法あるいはそれらの方法の組合せによって達成することができるが、本発明者らの研究では、通常の方法ではかなりの困難を伴う。一般に、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の製造においては、原料粗粉を微粉砕によって微粉化し、この微粉を磁界中で金型成形して成形体を得、これを焼結して焼結体とする方法が採られる。例えば、微粉砕をジェットミルを用いて行う場合には、粉砕時のガスの圧力や粗粉の供給速度等を制御することにより、所定の平均粒度や粒度分布を持つ微粉を得ることができる。また、必要に応じて、分級をおこなうことにより、微粉の粒度分布を制御することもできる。このようにして作製した微粉を成形し、焼結するにあたっては、さらに適切な焼結温度・時間・パターンを選択することによって、Ｒ−Ｆｅ−Ｂ系焼結型永久磁石の主相の結晶粒径を上記の規定範囲のものとすることは必ずしも不可能ではない。しかし、多くの条件を設定し、これを制御する必要があり、所定の結晶粒径を有する焼結体を再現性よく製造するのははなはだ困難であることが判った。
【００１３】
本発明者らは前記特定組成を有するＲ−Ｆｅ−Ｂ系焼結型永久磁石の主相の結晶粒径を上記の規定範囲とするのに容易で量産上適した方法を探索した結果、いわゆるストリップキャスト法と呼ばれる方法で製造された所定の組成を有するＲ−Ｆｅ−Ｂ系急冷薄帯状合金を、所定の温度範囲で熱処理し、これを粉砕して原料粗粉とする方法を見い出した。また熱処理後の薄帯状合金を粉砕するにあたっては、水素吸蔵により自然崩壊させた後脱水素処理を施してから行うことが微粉砕性能を高めるうえで有効である。
図６は、重量百分比率でＮｄ２７．８％，Ｐｒ０．４５％，Ｄｙ１．７％，Ｂ１．０５％，Ａｌ０．０５％，Ｃｏ２．０５％，Ｇａ０．０８％，Ｃｕ０．０９％，Ｏ０．０２％，Ｎ０．００４％，Ｃ０．００７％，残部Ｆｅの組成を有する、ストリップキャスト法で製造された薄帯状合金の断面組織を示す図である（ａｓｃａｓｔ）。デンドライト状の微細な組織が存在している。写真の中で白色に観察される相は希土類量が少なく永久磁石焼結体の主相に相当する相、黒色に観察される相は希土類量が多い永久磁石焼結体のＲｒｉｃｈ相に相当する相である。このＲｒｉｃｈ相は微粉砕時に破壊の起点となるので、このＲｒｉｃｈ相が図６に示すように微細に分散している薄帯状合金を使用した場合、粒径が細かくて均一な微粉が確率的に生成しやすい。従って、微粉砕時や焼結時の多くの条件を厳密に管理することなく、比較的容易にしかも再現性よく本発明の磁石主相粒径分布を有する焼結体が製造可能となるのである。
しかしこの薄帯状合金（急冷鋳造のまま）をこのまま直接粉砕して原料粗粉とし、これを微粉砕しても、良好な微粉の粒度分布を得るには困難を伴い、これを成形・焼結した焼結体では望ましい主相結晶粒径分布のものを再現性よく得ることが困難である。この理由は、急冷鋳造によって薄帯状合金の表面が硬化し、微粉砕時の被粉砕性をいちじるしく悪化させるからである。
【００１４】
本発明者らは、この問題を解決する手段として、この薄帯状合金を特定温度範囲で熱処理して薄帯状合金表面の硬化を除去することが有効であることを見い出した。熱処理の温度は８００℃〜１１００℃とされる。これは、熱処理温度が８００℃未満では硬化の除去が不十分だからである。また、１１００℃より高い温度では、熱処理時に薄帯状合金間で反応が生じ、後工程での処理に困難が生じるからである。活性な希土類元素を多量に含有する薄帯状合金であるため、熱処理は不活性ガス雰囲気中又は実質的な真空中で行う必要がある。また、前記のように、熱処理後の薄帯状合金に水素を吸蔵させて自然崩壊させ、脱水素処理をおこなった後、これを粗粉化することは、微粉砕性を高めるうえでさらに有効である。これは、熱処理による薄帯状合金表面の硬化の除去効果に加え、水素による薄帯状合金内部の主にはＲｒｉｃｈ相のぜい化効果が加わることによる。
【００１５】
表１に、図６の薄帯状合金を各種条件で熱処理（１Ｈｒ）あるいは粉砕して粗粉とし、これを同一条件で微粉砕し、成形・焼結した場合の焼結体の主相結晶粒径の状態を示す。なお、微粉砕以降の焼結体の製造方法・条件については、詳しく後述する。
【００１６】
【表１】

【００１７】
表１から、薄帯状合金を８００℃以上の温度で熱処理し、これを用いることによって、本発明の主相粒径分布を有する焼結体が得られることがわかる。また、前述したように、水素処理の有効性も明かである。同時に表１から、７００℃で熱処理したものの主相粒径の状態は、急冷鋳造したままのものとほぼ同水準である。よって、７００℃の熱処理温度では、薄帯状合金の表面硬化部の除去が不十分であることがわかる。
同時に本発明者らは、薄帯状合金に対して８００℃以上で熱処理をおこなうことにより、磁気特性のうち特にＢｒを向上できることを見い出した。結果を同じく表１に示す。表１から、急冷鋳造状態と７００℃で熱処理した薄帯状合金による永久磁石焼結体のＢｒは１３．２〜１３．３ＫＧである。これに対し、８００℃と９００℃で熱処理した薄帯状合金を使用した場合には、Ｂｒは１３．５５ＫＧと急激に増加する。熱処理温度が１０００℃では、結果として得られるＢｒは微増し、１３．６ＫＧとなる。１１００℃，１２００℃の熱処理温度では、Ｂｒの増加は飽和に達し、１３．６ＫＧと変わらない。
【００１８】
表１に示した薄帯状合金のうち、熱処理温度が９００℃，１０００℃，１１００℃のものの断面組織を、それぞれ図７、図８、図９に示す。急冷鋳造のままの状態（図６）をも含めこれらを対比すると、熱処理温度の上昇に従って、薄帯状合金内の主相に相当する白色組織、Ｒｒｉｃｈ相に相当する黒色組織のいずれもが粗大化していることがわかる。これらのことから本発明者らは、急冷鋳造のままの薄帯状合金では主相およびＲｒｉｃｈ相に相当する相から構成される組織が微細であるために、これを用いて微粉を製造した場合、微粉の内に多結晶状態のままのものが確率的に多く存在し、微粉を磁界中で金型成形する際の配向性の低下を招き、永久磁石焼結体のＢｒ低下をもたらしているものと考える。７００℃の熱処理温度では、上記組織の成長が不十分で、配向性の改善には至らない。上記図６〜図９に示すように、熱処理温度の上昇に従って薄帯状合金の内部組織が粗大化しているが、これによって多結晶状態の微粉の発生の確率が低下し、Ｂｒが改善されると考えられるが、表１の結果から判断する限り、８００℃の熱処理温度でその効果はかなり出ているものと考えられる。薄帯状合金の熱処理温度のさらなる増加に従って、得られる焼結体のＢｒはやや向上するものの、１０００℃以上の熱処理温度では飽和の傾向を示す。これは、薄帯状合金内部の組織がある程度粗大化し、多結晶状態の微粉が確率的にほとんど発生しない状態に達した段階では、熱処理温度をさらに上げて組織の粗大化を促進させても、それは得られる焼結体のＢｒの向上として反映しないということで理解できる。
【００１９】
以上詳細に説明したように、ストリップキャスト法による所定の組成の急冷鋳造薄帯状合金を、特定の温度範囲において熱処理し、あるいはこれに水素吸蔵処理を施して自然崩壊させ、これを粉砕して粗粉化することによって、微粉砕時の粉砕性が改善され、これを用いて製造された永久磁石焼結体は、耐蝕性にきわめて優れた本発明の主相結晶粒径を有するものとなるのであるが、同時に高い磁気特性を有するものにもなるのである。
なお、薄帯状合金の熱処理温度は８００〜１１００℃が好ましく、熱処理時間は好ましくは１５分間〜３時間、より好ましくは３０分間〜３時間である。
【００２０】
以下では、本発明のR-Fe-B系焼結型永久磁石の組成の限定理由を述べる。
希土類元素の量は、重量百分率で27.0〜31.0％とされる。希土類元素の量が31.0 ％を越えると、焼結体内部のRrich相の量が多くなり、かつ形態も粗大化して耐蝕性が悪くなる。一方、希土類元素の量が27.0％未満であると、焼結体の緻密化に必要な液相量が不足して焼結体密度が低下し、同時に磁気特性のうち残留磁束密度Brと保磁力iHcが共に低下する。従って、希土類元素の量は27.0〜31.0％とされる。
Oの量は重量百分率で0.05〜0.25％とされる。Oの量が0.25％を越える場合には、希土類元素の一部が酸化物を形成し、磁気的に有効な希土類元素が減少して保磁力iHcが低下する。一方溶解によって作製するインゴットのO量の水準は最大0.04％であるため、最終焼結体のO量をこの値以下とすることは困難であり、O量は0.05〜0.25％とすることが好ましい。
Cの量は重量百分率で0.01〜0.15％とされる。Cの量が0.15％より多い場合には、希土類元素の一部が炭化物を形成し、磁気的に有効な希土類元素が減少して保磁力iHcが低下する。C量は、0.12％以下とすることがより好ましく、0.10％以下とすることがさらに好ましい。一方、溶解によって作製するインゴットのC量の水準は最大0.008％であり、最終焼結体のC量をこの値以下とすることは困難であり、焼結体のC量は0.01〜0.15％とすることが好ましい。なお、焼結体のO量とC 量を上記値にする具体的な方法は後記する。
【００２１】
本発明者らの研究成果によると、Ｒ−Ｆｅ−Ｂ系焼結型希土類磁石の耐蝕性の大幅な改善に対しては、希土類元素の量を３１．０％以下とすることと先に述べた焼結体主相結晶粒径の大きさを前記特定範囲とすることは、必要条件ではあるが十分条件ではない。さらに、焼結体中のＮ量を厳密に制御する必要がある。前記特定範囲の焼結体主相結晶粒径を有し、上記の組成範囲の希土類量、Ｏ量、Ｃ量を有するＲ−Ｆｅ−Ｂ系焼結型希土類磁石において、焼結体中のＮ（窒素）量を所定範囲とすることによって、優れた耐蝕性と高い磁気特性を両立させることができる。焼結体中のＮ量は重量百分率で０．０２〜０．１５％とする必要がある。Ｎの含有による耐蝕性の改善効果のメカニズムは必ずしも明確ではないが焼結体中のＮは主にＲｒｉｃｈ相に存在し、希土類元素の一部と結合して窒化物を形成していることから、この窒化物の形成がＲｒｉｃｈ相の陽極酸化を抑制しているものと考えられる。Ｎの量が０．０２％より少ない場合には、窒化物の形成量が少ないためか、焼結体の耐蝕性の改善効果は見られない。Ｎの量が０．０２％以上では、Ｎの量の増加に従って焼結体の耐蝕性も向上するが、Ｎの量が０．１５％を越えると保磁力ｉＨｃが急激に低下する。これは、窒化物の形成による磁気的に有効な希土類元素の減少によるためと考えられる。以上の理由から、Ｎ量は０．０２〜０．１５％とされる。Ｎ量は０．０３〜０．１３％とすることがさらに好ましい。
【００２２】
本発明のR-Fe-B系焼結型永久磁石においては、Feの一部をNb,Al,Co,Ga,Cuのうちの１種又は２種以上で置換することができ以下に各元素の置換量（ここでは置換後の永久磁石の全組成に対する重量百分率）の限定の理由を説明する。
Nbの置換量は0.1〜2.0％とされる。Nbの添加によって、焼結過程でNbのほう化物が生成し、これが結晶粒の異常粒成長を抑制する。Nbの置換量が0.1％より少ない場合には、結晶粒の異常粒成長の抑制効果が十分ではなくなる。一方、Nbの置換量が2.0％を越えると、Nbのほう化物の生成量が多くなるため残留磁束密度Brが低下する。
Alの置換量は0.02〜2.0％とされる。Alの添加は保磁力iHcを高める効果がある。Alの置換量が0.02％より少ない場合には、保磁力の向上効果が少ない。置換量が2.0％を越えると、残留磁束密度Brが急激に低下する。
Coの置換量は0.3〜5.0％とされる。Coの添加はキュリー点の向上即ち飽和磁化の温度係数の改善をもたらす。Coの置換量が0.3％より少ない場合には、温度係数の改善効果は小さい。Coの置換量が5.0％を越えると、残留磁束密度Br、保磁力iHcが共に急激に低下する。
Gaの置換量は0.01〜0.5％とされる。Gaの微量添加は保磁力iHcの向上をもたらすが、置換量が0.01％より少ない場合には、添加効果は小さい。一方、Gaの置換量が0.5％を越えると、残留磁束密度Brの低下が顕著になるとともに保磁力iHcも低下する。
Cuの置換量は0.01〜1.0％とされる。Cuの微量添加は保磁力iHcの向上をもたらすが、置換量が1.0％を越えるとその添加効果は飽和する。添加量が0.01％より少ない場合には、保磁力iHcの向上効果は小さい。
【００２３】
次に、本発明の要点であるＲ−Ｆｅ−Ｂ系焼結型永久磁石のＮ量の制御方法について説明する。
Ｒ−Ｆｅ−Ｂ系焼結型永久磁石のＮ量の制御方法には種々の方法がありその方法は本発明においては選択可能であり、限定されるものではない。例えば、ジェットミル粉砕機にＲ−Ｆｅ−Ｂ系焼結型永久磁石用の原料粗粉を装入し、次いでジェットミル内部をＡｒガスで置換してそのＡｒガス中の酸素濃度が実質的に０％になるようにし、次にＮ_２ガスを微量導入してＡｒガス中のＮ_２ガスの濃度を調整する（通常０．０００１〜０．１ｖｏｌ％の範囲）。このＮ_２ガスを微量に含んだＡｒガス雰囲気中で原料粗粉を微粉砕する過程で、原料中の主には希土類元素とＮが結合し、回収された微粉中のＮ量が増加する。微粉の回収にあたっては、ジェットミルの微粉回収口に鉱物油、植物油、合成油等の溶媒を満たした容器を直接設置し、Ａｒガス雰囲気中で溶媒中に直接微粉を回収する。こうして得たスラリー状の原料を磁界中で湿式成形し、成形体とする。成形体を真空炉中で、５×１０^−２ｔｏｒｒ程度の真空度下で２００℃前後の温度に加熱し、成形体内の含有溶媒を除去する。次いで引き続き、真空炉の温度を１１００℃前後の焼結温度にまで上げ、５×１０^−４ｔｏｒｒ程度の真空度下で焼結して焼結体を得る。こうしてＯ量が０．２５％以下でＣ量が０．１５％以下のＲ−Ｆｅ−Ｂ系焼結型永久磁石を得ることができる。この場合、焼結体中のＮ量の制御は、上記粉砕時のＡｒガス中の導入Ｎ_２ガスの濃度制御によっておこなう。原料へのＮの混入度は、ジェットミルの容量、装入原料粗粉の組成と装入量、ジェットミル粉砕時の原料粗粉の送り量などによって変化する。
従って、目標とする焼結体Ｎ量を得るためには、粉砕時の条件毎に条件出しをし、最適なＡｒガス中のＮ_２ガス濃度を決めて粉砕する必要がある。この様な方法によって、焼結体中のＮ量を０．０２〜０．１５％に制御することができる。
【００２４】
また、ジェットミル内部をＮ_２ガスで置換してそのＮ_２ガス中の酸素濃度が実質的に０％になるようにし、このＮ_２ガス雰囲気中で原料粗粉を微粉砕することで、Ｏ量が０．２５％以下、Ｃ量が０．１５％以下、Ｎ量が０．０２〜０．１５％のＲ−Ｆｅ−Ｂ系焼結型永久磁石を得ることもできる。この場合は、原料粗粉の装入量と粉砕時の原料粗粉の送り量によって原料へのＮの混入度を制御し、目標とするＮ量の焼結体を得る。ジェットミルの型式や容量によって原料へのＮの混入度は変化するため、あらかじめ条件出しを行って、原料粗粉の装入量と粉砕時の送り量を設定する。粉砕後の微粉の回収方法は鉱物油、植物油、合成油のうちの１種又は２種以上からなる溶媒中へであり、湿式成形以降の工程も前記のＡｒガス雰囲気中での粉砕の場合と同じである。
なお、以上に述べた酸素濃度が実質的に０％である雰囲気とは、例えばＲ−Ｆｅ−Ｂ系原料粗粉を１０ｋｇ／Ｈｒ程度微粉砕できる能力を有する生産型のジットミル粉砕機の場合では、雰囲気中の酸素濃度が百分比率で０．０１ｖｏｌ％以下、より好ましくは０．００５ｖｏｌ％以下、さらに好ましくは０．００２ｖｏｌ％以下の雰囲気を言う。
以上のような方法によってＯ量が０．２５％以下、Ｃ量が０．１５％以下、Ｎ量が０．０２〜０．１５％のＲ−Ｆｅ−Ｂ系焼結体をつくることができるが、同時に、先に説明した８００〜１１００℃の温度範囲で熱処理を施した所定の組成を有する急冷薄帯状合金を原料として用いることによって、前記特定範囲の主相結晶粒径のものが容易にかつ再現性よく得られる。
こうして得られた焼結体を熱処理、加工することによって、耐蝕性に優れかつ高い磁気特性を有するＲ−Ｆｅ−Ｂ系焼結型永久磁石の製造が可能である。
【００２５】
【発明の実施の形態】
以下、本発明を実施例をもって具体的に説明するが、本発明の内容はこれに限定されるものではない。
（実施例１）
重量百分率でＮｄ２７．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．３５％，Ａｌ０．０８％，Ｃｏ２．５％，Ｇａ０．０９％，Ｃｕ０．０８％，Ｏ０．０３％，Ｃ０．００５％，Ｎ０．００４％，残部Ｆｅの組成を有する、厚さが０．２〜０．５ｍｍの薄帯状合金を、ストリップキャスト法で作製した。この薄帯状の合金を、Ａｒガス雰囲気中で１０００℃で２時間加熱した。次に水素炉を使用し、この薄帯状の合金を常温で水素ガス雰囲気中で水素吸蔵させ、自然崩壊させた。次いで炉内を真空排気しつつ５５０℃まで薄帯状の合金を加熱し、その温度で１時間保持して脱水素処理を行った。崩壊した合金を窒素ガス雰囲気中で機械的に破砕して、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、Ｎｄ２７．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．３５％，Ａｌ０．０８％，Ｃｏ２．５％，Ｇａ０．０９％，Ｃｕ０．０８％，Ｏ０．１２％，Ｃ０．０２％，Ｎ０．００８％，残部Ｆｅという分析値を得た。
この原料粗粉８０ｋｇをジェットミル内に装入した後、ジェットミル内部をＮ_２ガスで置換し、Ｎ_２ガス中の酸素濃度を実質的に０％（酸素分析計値で０．００１ｖｏｌ％）とした。次いで、粉砕圧力７．０ｋｇ／ｃｍ^２、原料粗粉の供給量１０ｋｇ／Ｈｒの条件で粉砕した。微粉の平均粒度は３．９μｍであった。
ジェットミルの微粉回収口には鉱物油（商品名出光スーパーゾルＰＡ−３０，出光興産製）を満たした容器を直接設置し、Ｎ_２ガス雰囲気中で微粉を直接鉱物油中へ回収した。回収後の原料は、鉱物油の量を加減することで微粉の純分が８０重量％の原料スラリーとした。この原料スラリーを、金型キャビティ内で１２ｋＯｅの配向磁界を印加しながら０．８ｔｏｎ／ｃｍ^２の成形圧で湿式成形した。配向磁界の印加方向は、成形方向と垂直である。また、金型の上パンチには溶媒排出孔を多数設け、成形時には１ｍｍの厚さの布製のフィルタを上パンチ面にあてて使用した。
成形体は、５．０×１０^−２ｔｏｒｒの真空中で２００℃×１時間加熱して含有鉱物油を除去し、次いで４．０×１０^−４ｔｏｒｒの条件下で１５℃／分の昇温速度で１０７０℃まで昇温し、その温度で３時間保持して焼結した。
焼結体の組成を分析したところ、Ｎｄ２７．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．３５％，Ａｌ０．０８％，Ｃｏ２．５％，Ｇａ０．０９％，Ｃｕ０．０８％，Ｏ０．１６％，Ｃ０．０７％，Ｎ０．０５５％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は９３％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は４％であった。
この焼結体にＡｒガス雰囲気中で９００℃×２時間と４８０℃×１時間の熱処理を各１回施した。機械加工後磁気特性を測定したところ、表２に示すような良好な値を得た。
この永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの一定寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施した。次いでこの試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べた。表２に示すように、２５００時間を経過してもＮｉメッキに異常が認められず、良好な耐蝕性を示した。
【００２６】
（実施例２）
重量百分率でＮｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．０２％，Ｃ０．００５％，Ｎ０．００３％，残部Ｆｅの組成を有する、厚さが０．２〜０．４ｍｍの薄帯状合金を、ストリップキャスト法で作製した。この薄帯状の合金を、Ａｒガス雰囲気中で１１００℃で１時間加熱した。次に水素炉を使用し、この薄帯状の合金を常温で水素ガス雰囲気中で水素吸蔵させ、自然崩壊させた。次いで炉内を真空排気しつつ５５０℃まで薄帯状の合金を加熱し、その温度で１時間保持して脱水素処理を行った。崩壊した合金を窒素ガス雰囲気中で機械的に破砕して、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、Ｎｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１１％，Ｃ０．０２％，Ｎ０．００６％，残部Ｆｅという分析値を得た。
この原料粗粉１００ｋｇをジェットミル内に装入した後、ジェットミル内部をＮ_２ガスで置換し、Ｎ_２ガス中の酸素濃度を実質的に０％（酸素分析計値で０．００２％）とした
。次いで、粉砕圧力８．０ｋｇ／ｃｍ^２、原料粗粉の供給量１２ｋｇ／Ｈｒの条件で粉砕した。微粉の平均粒度は３．８μｍであった。
ジェットミルの微粉回収口には鉱物油（商品名出光スーパーゾルＰＡ−３０，出光興産製）を満たした容器を直接設置し、Ｎ_２ガス雰囲気中で微粉を直接鉱物油中へ回収した。回収後の原料は、鉱物油の量を加減することで微粉の純分が７７重量％の原料スラリーとした。この原料スラリーを、金型キャビティ内で１０ｋＯｅの配向磁界を印加しながら１．５ｔｏｎ／ｃｍ^２の成形圧で湿式成形した。配向磁界の印加方向は、成形方向と垂直である。また、金型の上パンチには溶媒排出孔を多数設け、成形時には１ｍｍの厚さの布製のフィルタを上パンチ面にあてて使用した。
成形体は、５．０×１０^−２ｔｏｒｒの真空中で２００℃×２時間加熱して含有鉱物油を除去し、次いで５．０×１０^−４ｔｏｒｒの条件下で１５℃／分の昇温速度で１０９０℃まで昇温し、その温度で３時間保持して焼結した。
焼結体の組成を分析したところ、Ｎｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１４％，Ｃ０．０６％，Ｎ０．０４０％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は９５％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は３％であった。この焼結体の粒径と面積率との関係を図１０に示す。
この焼結体にＡｒガス雰囲気中で９００℃×２時間と４６０℃×１時間の熱処理を各１回施した。機械加工後磁気特性を測定したところ、表２に示すような良好な値を得た。
この永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの一定寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施した。次いでこの試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べた。表２に示すように、２５００時間を経過してもＮｉメッキに異常が認められず、良好な耐蝕性を示した。
また、図１０より、本発明にかかる焼結体は結晶粒径５μｍを中心にシャープな粒径分布となっていることがわかる。
【００２７】
（実施例３）
重量百分率でＮｄ２０．７％，Ｐｒ８．６％，Ｄｙ１．２％，Ｂ１．０５％，Ａｌ０．０８％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．０３％，Ｃ０．００６％，Ｎ０．００４％，残部Ｆｅの組成を有する、厚さが０．１〜０．５ｍｍの薄帯状合金を、ストリップキャスト法で作製した。この薄帯状の合金を、Ａｒガス雰囲気中で９００℃で３時間加熱した。次に水素炉を使用し、この薄帯状の合金を常温で水素ガス雰囲気中で水素吸蔵させ、自然崩壊させた。次いで炉内を真空排気しつつ５５０℃まで薄帯状の合金を加熱し、その温度で１時間保持して脱水素処理を行った。崩壊した合金を窒素ガス雰囲気中で機械的に破砕して、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、Ｎｄ２０．７％，Ｐｒ８．６％，Ｄｙ１．５％，Ｂ１．０５％，Ａｌ０．０８％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１３％，Ｃ０．０３％，Ｎ０．００９％，残部Ｆｅという分析値を得た。
この原料粗粉５０ｋｇをジェットミル内に装入した後、ジェットミル内部をＡｒガスで置換し、Ａｒガス中の酸素濃度を実質的に０％（酸素分析計値で０．００２ｖｏｌ％）とした。次にＡｒガス中のＮ_２ガスの濃度を０．００５ｖｏｌ％とした。次いで、粉砕圧力７．５ｋｇ／ｃｍ^２、原料粗粉の供給量８ｋｇ／Ｈｒの条件で粉砕した。
ジェットミルの微粉回収口には鉱物油（商品名出光スーパーゾルＰＡ−３０，出光興産製）を満たした容器を直接設置し、Ａｒガス雰囲気中で微粉を直接鉱物油中へ回収した。回収後の原料は、鉱物油の量を加減することで微粉の純分が７５重量％の原料スラリーとした。なお、微粉の平均粒度は４．０μｍであった。この原料スラリーを、金型キャビティ内で１３ｋＯｅの配向磁界を印加しながら０．６ｔｏｎ／ｃｍ^２の成形圧で湿式成形した。配向磁界の印加方向は、成形方向と垂直である。また、金型の上パンチには溶媒排出孔を多数設け、成形時には１ｍｍの厚さの布製のフィルタを上パンチ面にあてて使用した。
成形体は、６．０×１０^−２ｔｏｒｒの真空中で１８０℃×４時間加熱して含有鉱物油を除去し、次いで３．０×１０^−４ｔｏｒｒの条件下で１５℃／分の昇温速度で１０７０℃まで昇温し、その温度で２時間保持して焼結した。
焼結体の組成を分析したところ、Ｎｄ２０．７％，Ｐｒ８．６％，Ｄｙ１．２％，Ｂ１．０５％，Ａｌ０．０８％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１８％，Ｃ０．０７％，Ｎ０．０７５％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は９０％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は５％であった。
この焼結体にＡｒガス雰囲気中で９００℃×２時間と５１０℃×１時間の熱処理を各１回施した。機械加工後磁気特性を測定したところ、表２に示すような良好な値を得た。
この永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの一定寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施した。次いでこの試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べた。表２に示すように、２５００時間を経過してもＮｉメッキに異常が認められず、良好な耐蝕性を示した。
【００２８】
（実施例４）
重量百分率でＮｄ２２．０％，Ｐｒ５．０％，Ｄｙ１．５％，Ｂ１．１％，Ａｌ１．０％，Ｃｏ２．５％，Ｏ０．０２％，Ｃ０．００５％，Ｎ０．００５％，残部Ｆｅの組成を有する、厚さが０．１〜０．４ｍｍの薄帯状合金を、ストリップキャスト法で作製した。この薄帯状の合金を、Ａｒガス雰囲気中で１０００℃で２時間加熱した。熱処理後の薄帯状合金を窒素ガス雰囲気中で機械的に破砕して、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、Ｎｄ２２．０％，Ｐｒ５．０％，Ｄｙ１．５％，Ｂ１．１％，Ａｌ１．０％，Ｃｏ２．５％，Ｏ０．１４％，Ｃ０．０１％，Ｎ０．００９％，残部Ｆｅという分析値を得た。この原料粗粉５０ｋｇをジェットミル内に装入した後、ジェットミル内部をＮ_２ガスで置換し、Ｎ_２ガス中の酸素濃度を実質的に０％（酸素分析計値で０．００２ｖｏｌ％）とした。次いで、粉砕圧力７．０ｋｇ／ｃｍ^２、原料粗粉の供給量１０ｋｇ／Ｈｒの条件で粉砕した。微粉の平均粒度は４．２μｍであった。
ジェットミルの微粉回収口には鉱物油（商品名出光スーパーゾルＰＡ−３０，出光興産製）を満たした容器を直接設置し、Ｎ_２ガス雰囲気中で微粉を直接鉱物油中へ回収した。回収後の原料は、鉱物油の量を加減することで微粉の純分が７８重量％の原料スラリーとした。この原料スラリーを、金型キャビティ内で１１ｋＯｅの配向磁界を印加しながら０．５ｔｏｎ／ｃｍ^２の成形圧で湿式成形した。配向磁界の印加方向は、成形方向と垂直である。また金型の上パンチには溶媒排出孔を多数設け、成形時には１ｍｍの厚さの布製のフィルタを上パンチ面にあてて使用した。
成形体は、５．０×１０^−２ｔｏｒｒの真空中で２００℃×２時間加熱して含有鉱物油を除去し、次いで２．０×１０^−４ｔｏｒｒの条件下で１５℃／分の昇温速度で１０８０℃まで昇温し、その温度で３時間保持して焼結した。
焼結体の組成を分析したところ、Ｎｄ２２．０％，Ｐｒ５．０％，Ｄｙ１．５％，Ｂ１．１％，Ａｌ１．０％，Ｃｏ２．５％，Ｏ０．１７％，Ｃ０．０７％，Ｎ０．０６０％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は８８％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は７％であった。
この焼結体にＡｒガス雰囲気中で９００℃×２時間と６００℃×１時間の熱処理を各１回施した。機械加工後磁気特性を測定したところ、表２に示すような良好な値を得た。
この永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの一定寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施した。次いでこの試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べた。表２に示すように、２０００時間を経過してもＮｉメッキに異常が認められず、良好な耐蝕性を示した。
【００２９】
（比較例１）
実施例１で作製した薄帯状の合金を、熱処理をおこなわずに直接水素炉に入れ、常温で水素ガス雰囲気中で水素吸蔵させ、自然崩壊させた。その後、実施例１と同じ条件で脱水素処理と機械的破砕をおこない、３２ｍｅｓｈ以下の原料粗粉とした。この原料粗粉の組成を分析したところ、重量百分率でＮｄ２７．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．３５％，Ａｌ０．０８％，Ｃｏ２．５％，Ｇａ０．０９％，Ｃｕ０．０８％，Ｏ０．１０％，Ｃ０．０２％，Ｎ０．００７％，残部Ｆｅという分析値を得た。
この原料粗粉を、実施例１と同一の条件で微粉砕した。得られた微粉の平均粒度は４．４μｍと、実施例１の場合に比べて粗かった。
微粉の回収、原料スラリーの作製、湿式成形、脱鉱物油と焼結、熱処理、耐蝕性の評価などの以降の工程も、実施例１と同一の条件でおこなった。焼結体の組成を分析したところ、Ｎｄ２７．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．３５％，Ａｌ０．０８％，Ｃｏ２．５％，Ｇａ０．０９％，Ｃｕ０．０８％，Ｏ０．１４％，Ｃ０．０６％，Ｎ０．０４５％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は７８％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は１２％であった。
この永久磁石の磁気特性を評価したところ、表２に示すように、実施例１の値に比べてＢｒ，ｉＨｃ共若干低い値であった。また、この永久磁石の耐蝕性は、表２に示すように１２００時間を経過してもＮｉメッキに異常が認められず実用上全く問題ない水準にあることがわかったが、２０００時間の経過でＮｉメッキのわずかなハク離が発生し、実施例１で製造した焼結体との比較では耐蝕性に劣ることが判明した。
【００３０】
（比較例２）
実施例２と同一の組成を有するＲ−Ｆｅ−Ｂ系合金インゴットを作製した。この合金の組成分析値は重量百分比率でＮｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．０１％，Ｃ０．００４％，Ｎ０．００２％，残部Ｆｅであった。合金の組織中にα−Ｆｅの析出が認められたため、これを消去するため、合金インゴットにアルゴンガス雰囲気中で１１００℃×６時間の液体化処理を施した。次に合金インゴットを水素炉中に入れ、常温で水素吸蔵させて自然崩壊させた。自然崩壊後の合金に、実施例２と同一の条件で脱水素処理と機械的破砕とを行い、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、重量百分率でＮｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１０％，Ｃ０．０２％，Ｎ０．００５％，残部Ｆｅという分析値を得た。
この原料粗粉を、実施例２と同一の条件で微粉砕した。得られた微粉の平均粒度は４．７μｍと、実施例１の場合に比べて粗かった。
微粉の回収、原料スラリーの作製、湿式成形、脱鉱物油と焼結、熱処理、耐蝕性の評価などの以降の工程も、実施例２と同一の条件でおこなった。焼結体の組成を分析したところ、Ｎｄ２２．３％，Ｐｒ２．０％，Ｄｙ５．５％，Ｂ１．０％，Ｎｂ０．５％，Ａｌ０．２％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１２％，Ｃ０．０６％，Ｎ０．０３０％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は６１％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は２２％であった。この焼結体の粒径と面積率との関係を図１１に示す。
この永久磁石の磁気特性を評価したところ、表２に示すように、実施例２の値とほぼ同等の良好な値であった。また、この永久磁石の耐蝕性は、表２に示すように１０００時間を経過してもＮｉメッキに異常が認められず実用上全く問題ない水準にあることがわかったが、１９００時間の経過でＮｉメッキの一部にわずかなハク離が発生し、実施例２で製造した永久磁石との比較では耐蝕性に劣ることが判明した。また、図１１より、比較例２の焼結体は結晶粒径８μｍを中心にブロードな粒径分布となっており、実施例２の焼結体の粒径分布である図１０と比較すると１３μｍ以上の結晶粒径の面積率が多いことがわかる。
【００３１】
（比較例３）
実施例３で使用したのと同一の原料粗粉を、実施例３と同一の条件で微粉砕した。ただしＡｒガス中にＮ_２ガスは導入しなかった。微粉の平均粒度は４．０μｍであった。微粉の回収、原料スラリーの作製、湿式成形、脱鉱物油と焼結、熱処理、耐蝕性の評価などの以降の工程も、実施例３と同一の条件でおこなった。焼結体の組成を分析したところ、Ｎｄ２０．７％，Ｐｒ８．６％，Ｄｙ１．２％，Ｂ１．０５％，Ａｌ０．０８％，Ｃｏ２．０％，Ｇａ０．０９％，Ｃｕ０．１％，Ｏ０．１８％，Ｃ０．０７％，Ｎ０．０１０％，残部Ｆｅという分析値を得た。
この焼結体の磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は９２％、１３μｍ以上の主相結晶粒の面積の和は４％であった。
この永久磁石の磁気特性を評価したところ、表１に示すような良好な値を得た。しかし、この永久磁石の耐蝕性は、表２に示すように１９２時間でＮｉメッキにハク離が発生し、良好なものではないことが判った。
【００３２】
（比較例４）
重量百分率でＮｄ３０．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．８％，Ａｌ０．２％，Ｃｏ３．０％，Ｇａ０．０８％，Ｃｕ０．１％，Ｏ０．０２％，Ｃ０．００５％，Ｎ０．００５％，残部Ｆｅの組成を有する、厚さが０．２〜０．５ｍｍの薄帯状合金を、ストリップキャスト法で作製した。この薄帯状の合金を、Ａｒガス雰囲気中で９５０℃で４時間加熱した。次に水素炉を使用し、この薄帯状の合金を常温で水素ガス雰囲気中で水素吸蔵させ、自然崩壊させた。次いで炉内を真空排気しつつ５５０℃まで薄帯状の合金を加熱し、その温度で１時間保持して脱水素処理を行った。崩壊した合金を窒素ガス雰囲気中で機械的に破砕して、３２ｍｅｓｈ以下の原料粗粉とした。
この原料粗粉の組成を分析したところ、Ｎｄ３０．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．８％，Ａｌ０．２％，Ｃｏ３．０％，Ｇａ０．０８％，Ｃｕ０．１％，Ｏ０．１２％，Ｃ０．０２％，Ｎ０．００９％，残部Ｆｅという分析値を得た。
この原料粗粉１００ｋｇをジェットミル内に装入した後、ジェットミル内部をＮ_２ガスで置換し、Ｎ_２ガス中の酸素濃度を実質的に０％（酸素分析計値で０．００１ｖｏｌ％）とした。次いで、粉砕圧力７．５ｋｇ／ｃｍ^２、原料粗粉の供給量１０ｋｇ／Ｈｒの条件で粉砕した。微粉の平均粒度は４．１μｍであった。
ジェットミルの微粉回収口には鉱物油（商品名出光スーパーゾルＰＡ−３０，出光興産製）を満たした容器を直接設置し、Ｎ_２ガス雰囲気中で微粉を直接鉱物油中へ回収した。回収後の原料は、鉱物油の量を加減することで微粉の純分が７０重量％の原料スラリーとした。この原料スラリーを、金型キャビティ内で１４ｋＯｅの配向磁界を印加しながら０．８ｔｏｎ／ｃｍ^２の成形圧で湿式成形した。配向磁界の印加方向は、成形方向と垂直である。また金型の上パンチには溶媒排出孔を多数設け、成形時には１ｍｍの厚さの布製のフィルタを上パンチ面にあてて使用した。
成形体は、５．０×１０^−２ｔｏｒｒの真空中で２００℃×２時間加熱して含有鉱物油を除去し、次いで３．０×１０^−４ｔｏｒｒの条件下で１５℃／分の昇温速度で１０８０℃まで昇温し、その温度で３時間保持して焼結した。
焼結体の組成を分析したところ、Ｎｄ３０．０％，Ｐｒ０．５％，Ｄｙ１．５％，Ｂ１．０５％，Ｎｂ０．８％，Ａｌ０．２％，Ｃｏ３．０％，Ｇａ０．０８％，Ｃｕ０．１％，Ｏ０．１５％，Ｃ０．０７％，Ｎ０．０６０％，残部Ｆｅという分析値を得た。
この焼結体の、磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和は９２％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和は４％であった。
この焼結体にＡｒガス雰囲気中で９００℃×２時間と５５０℃×１時間の熱処理を各１回施した。機械加工後磁気特性を測定したところ、表２に示すような良好な値を得た。
この永久磁石の耐蝕性を評価するために、磁石を８ｍｍ×８ｍｍ×２ｍｍの一定寸法に加工後、その表面に膜厚１０μｍのＮｉメッキを施した。次いでこの試料を２気圧，１２０℃，湿度１００％の条件に放置し、時間の経過に対するＮｉメッキのハク離程度を調べたところ、表２に示すように４８時間でＮｉメッキにハク離が発生し、実用上全く適さないものであることが判った。
【００３３】
【表２】

【００３４】
【発明の効果】
本発明により、磁気特性を低下させずに、優れた耐食性を有する R-Fe-B 系焼結型永久磁石が得られる。
【図面の簡単な説明】
【図１】磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が９６％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１％である焼結型永久磁石の金属組織を示す図である。
【図２】磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が６４％，結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１７％の焼結型永久磁石の金属組織を示す図である。
【図３】磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が９６％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１％である焼結型永久磁石の耐食性評価試験５０００時間経過後の金属組織を示す図である。
【図４】磁石主相の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和が６４％，結晶粒径が１３μｍ以上の主相結晶粒の面積の和が１７％の焼結型永久磁石の耐食性評価試験２０００時間経過後の金属組織を示す図である。
【図５】図５は、磁石主相結晶の総面積に対する結晶粒径が１０μｍ以下の主相結晶粒の面積の和の割合と、磁石主相結晶の総面積に対する結晶粒径が１３μｍ以上の主相の結晶粒の面積の和の割合と、耐蝕性の加速試験での、Ｎｉメッキのハク離開始が生じるまでの経過時間との関係を示した図である。
【図６】ストリップキャスト法で作製した薄帯状合金の断面の金属組織を示す図である。
【図７】ストリップキャスト法で作製した薄帯状合金を９００℃で熱処理した後の断面の金属組織を示す図である。
【図８】ストリップキャスト法で作製した薄帯状合金を１０００℃で熱処理した後の断面の金属組織を示す図である。
【図９】ストリップキャスト法で作製した薄帯状合金を１１００℃で熱処理した後の断面の金属組織を示す図である。
【図１０】磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和が９５％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和が３％である焼結体の粒径分布を示す図である。
【図１１】磁石主相結晶の総面積に対する、結晶粒径が１０μｍ以下の主相結晶粒の面積の和が６１％、結晶粒径が１３μｍ以上の主相結晶粒の面積の和が２２％である焼結体の粒径分布を示す図である。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to improving the performance of sintered permanent magnets based on R-Fe-B.To.
[0002]
[Prior art]
Among sintered rare earth permanent magnets, R-Fe-B-based (R is one or more of rare earth elements including Y) sintered permanent magnets are attracting attention as high performance magnets and are used in a wide range of fields. Have been.
This R-Fe-B based sintered permanent magnet basically has an R₂Fe₁₄B phase (main phase), RFe₇B₆Phase (Brich phase), R₈₅Fe_FifteenIt has a structure composed of three phases (Rrich phase). Due to the existence of the Rrich phase rich in rare earth elements in composition and such a three-phase structure, the R-Fe-B-based sintered permanent magnet is smaller than the Sm-Co-based sintered permanent magnet. Poor corrosion resistance is one of the drawbacks from the beginning of the development of this permanent magnet to the present.
Although there is no established theory on the corrosion mechanism of the R-Fe-B based sintered permanent magnet, since the form of the corrosion starting from the Rrich phase is common, the view of anode corrosion using the Rrich phase as the anode is considered. There is also. Certainly, by reducing the amount of the rare earth element in the R-Fe-B based sintered permanent magnet, the amount of the Rrich phase inside the sintered body is reduced, and the morphology of the phase becomes finer. Correspondingly, the corrosion resistance of the permanent magnet is improved. Therefore, reducing the amount of the rare earth element is one method for improving the corrosion resistance of the R—Fe—B based sintered permanent magnet.
[0003]
A sintered rare earth permanent magnet containing an R-Fe-B system is a powder metallurgy process in which a raw metal is melted and poured into a mold, and the obtained ingot is pulverized, molded, sintered, heat-treated, and processed. It is generally manufactured by However, the alloy powder obtained by pulverizing the ingot is chemically very active because it contains a large amount of rare earth elements, and oxidizes in the atmosphere to increase the oxygen content. As a result, a part of the rare earth element of the obtained sintered body forms an oxide, and the magnetically effective rare earth element is reduced. For this reason, in order to realize a practical level of magnetic properties, for example, iHc ≧ 13 kOe, it is necessary to increase the amount of the rare earth element in the R—Fe—B based sintered permanent magnet, and the percentage by weight is 31%. The amount of rare earth element added exceeds that in practical materials.
For this reason, the corrosion resistance of the conventional R-Fe-B-based sintered permanent magnet was not sufficient.
[0004]
[Problems to be solved by the invention]
The present invention significantly improves the corrosion resistance of the R-Fe-B based sintered permanent magnets described above.Trying toTo do.
[0005]
[SectionMeans for solving the problem]
The present inventors have made various studies to improve the corrosion resistance of R-Fe-B based sintered permanent magnets, and as a result, a specific range of rare earth content and a specific amount of oxygen and carbon content of R-Fe- Nitrogen content within a specified range for B-based sintered permanent magnetsamountBy doing so, it has been found that corrosion resistance is improved and practically high magnetic properties are obtained. As a result of continuing research to further improve the corrosion resistance of R-Fe-B based sintered permanent magnets, the corrosion resistance was further improved by reducing the magnet main phase crystal grain size to a specified value or less. Have led to the present invention.
[0006]
Hereinafter, the present invention will be described specifically.
The present inventionSintered permanent magnetIs R in weight percentage (R is one or more of rare earth elements including Y) 27.0-31.0%, B 0.5-2.0%, N 0.02-0.15%, O 0.25% or less, C 0.15% or less , A sintered permanent magnet having a composition of the balance of Fe, wherein the R-Fe for the sintered permanent magnet has a two-phase structure substantially corresponding to a main phase and an Rrich phase of the sintered permanent magnet. Manufactured using -B quenched alloyCharacterized by improved corrosion resistance.
The present inventionSintered permanent magnetIs R in weight percentage (R is one or more of rare earth elements including Y) 27.0-31.0%, B 0.5-2.0%, N 0.02-0.15%, O 0.25% or less, C 0.15% or less A main phase crystal having a composition of the balance Fe of 80% or more of the area of the main phase crystal grains having a crystal grain size of 10 μm or less and a crystal grain size of 13 μm or more based on the total area of the magnet main phase crystal grains. A sintered permanent magnet having a total grain area of 10% or less for the sintered permanent magnet having a two-phase structure substantially corresponding to a main phase and an Rrich phase of the sintered permanent magnet. Manufactured using R-Fe-B quenched alloyCharacterized by improved corrosion resistance.
An R-Fe-B-based quenched alloy for the sintered permanent magnet having a two-phase structure substantially corresponding to the main phase and the Rrich phase of the sintered permanent magnet is a plate manufactured by a so-called strip casting method. It is a quenched thin strip alloy having a thickness of 0.1 to 0.5 mm and contains a small amount of a Brich phase in addition to the phase corresponding to the main phase and the Rrich phase of the sintered permanent magnet of the present invention. It consists of two phases, a phase corresponding to the phase and a phase corresponding to the Rrich phase.
In the sintered permanent magnet, a part of Fe is Nb 0.1 to 2.0%, Al 0.02 to 2.0%, Co 0.3 to 5.0%, Ga 0.01 to 0.5%, and Cu 0.01 to 1.0%.ofOne or two or more can be substituted.
[0007]
The present inventors have found that the corrosion resistance of the R-Fe-B based sintered permanent magnet having the above composition has a crystal grain size dependency, and the magnet main phase crystal grain size is particularly excellent by setting the crystal grain size to a specific value or less. It was found that corrosion resistance was exhibited. There are various methods for defining and measuring the magnet crystal grain size, which are not unique.However, the present inventors have determined that the sum of the area of the main phase crystal grains whose grain size is equal to or less than a certain dimension with respect to the total area of the magnet main phase. And the ratio of the sum of the area of the main phase crystal grains having a certain size or more to the total area of the magnet main phase as a scale of the magnet crystal grain size. Hereinafter, the effect of the present invention will be described using this scale. In measuring the ratio, the crystal structure of the target R-Fe-B-based sintered permanent magnet was observed with a microscope manufactured by OLYMPUS (trade name: VANOX), and this image was manufactured by NIRECO. This was performed by directly charging the image processing apparatus (product name LUZEX2).
[0008]
FIG. 1 shows Nd 27.5%, Pr 0.5%, Dy 1.5%, B 1.1%, Al 0.1%, Co 2.0%, Ga 0.08%, N 0.06% of O, 0.16% of O, 0.06% of C, 0.040% of N, and the balance of Fe, and the crystal grains of the main phase having a grain size of 10 μm or less with respect to the total area of the magnet main phase. It is an observation result by an optical microscope (observed at 1000 times) of a sintered permanent magnet having a total area of 96% and a total area of main phase crystal grains having a crystal grain size of 13 μm or more of 1%. FIG. 2 shows the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less with respect to the total area of the magnet main phase having the same composition of 64%, and the area of the main phase crystal grains having a crystal grain size of 13 μm or more. Is a result of observation by a light microscope (observed at 1000 times) of a sintered permanent magnet having a sum of 17%.
[0009]
In order to evaluate the corrosion resistance of these sintered permanent magnets, the magnets were processed into dimensions of 8 mm × 8 mm × 2 mm, and their surfaces were plated with Ni having a thickness of 20 μm. Next, the sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity, and the degree of separation of Ni plating over time was examined. In the sample produced from the permanent magnet of FIG. 1 having a small crystal grain size, no abnormality was observed in the Ni plating even after 2500 hours. With the passage of 5000 Hr, slight separation of Ni plating was observed. On the other hand, in the sample prepared from the permanent magnet of FIG. 2 having relatively large crystal grains, no abnormality was observed in the Ni plating even after 1000 hours. Since the above evaluation method is an acceleration test, it can be said from this result that the corrosion resistance of the permanent magnet in FIG. However, a large separation of Ni plating was observed in the sample after lapse of 2,000 hours, which indicates that there is a strict difference in corrosion resistance between the permanent magnet in FIG. 1 and the permanent magnet in FIG. Was. That is, the smaller the crystal grain size of the magnet main phase, the better the corrosion resistance. FIG. 3 is an SEM observation result of a cross section of the sample manufactured from the permanent magnet of FIG. 1 after the evaluation test was performed for 5000 hours. Although there is partial separation between the Ni plating and the permanent magnet sintered body as the base, the adhesion between the two is relatively good. Also, it can be seen from the 5000 Hr acceleration test that the permanent magnet sintered body as the base is hardly damaged.
FIG. 4 is an SEM observation result of a cross section of the sample manufactured from the sample of FIG. 2 after the above evaluation test was performed for 2000 hours. It can be seen from the accelerated test that the crystal grain boundaries of the permanent magnet sintered body, which is the base, were destroyed, thereby causing large separation of the Ni plating.
[0010]
From the above results, it was found that there is a difference in how the grain boundary is broken in the accelerated corrosion resistance test due to the difference in the main phase crystal grain size of the permanent magnet sintered body. Assuming the cause from FIG. 4, in the permanent magnet sintered body as shown in FIG. 2 in which relatively large main phase crystal grains are present, the voids between the main phase crystal grains, specifically, the grains are relatively small. The field triple point is the main part, in which the Ndrich phase, which is extremely susceptible to oxidation, is present, but the volume of the void filled with this Ndrich phase is large. Factors causing corrosion destruction, for example, water in the present accelerated test, are considered to be in a state in which the penetration of such factors is good and destruction of crystal grain boundaries is likely to occur in a chain reaction. The above is the explanation of the fact that the corrosion resistance of the sintered R-Fe-B permanent magnet according to the present invention depends on the crystal grain size of the main phase by showing an example of the research results of the present inventors. It is.
[0011]
The present inventors have continued the above-mentioned evaluation on the correlation between the composition, main phase crystal grain size distribution, magnetic properties, and corrosion resistance of the R-Fe-B based sintered permanent magnet employed in the present invention. And the results as shown in FIG.
FIG. 5 shows the ratio of the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less to the total area of the magnet main phase crystals, and the main phase having the crystal grain size of 13 μm or more with respect to the total area of the magnet main phase crystals. 2 shows the relationship between the ratio of the sum of the areas of the crystal grains and the elapsed time until the start of peeling of the Ni plating in the accelerated corrosion resistance test. The circles indicate Nd 25.5%, Pr 2.5%, Dy 2.0%, B 1.0%, Nb 0.4%, Al 1.0%, Co 3.0%, Cu 3.0% by weight percentage. A sintered body having a composition of 0.1%, 0.19% of O, 0.08% of C, 0.040% of N, and the balance of Fe, and □ indicates 28.0% of Nd and 1.0% of Dy in weight percentage. %, B 1.05%, Al 0.05%, Co 2.0%, Ga 0.09%, O 0.20%, C 0.07%, N 0.080%, and the balance Fe. The sintered body, and the triangles indicate Nd 24.5%, Pr 1.5%, Dy 4.5%, B 1.1%, Nb 1.0%, Al 0.2%, Co2. This shows a sintered body having a composition of 0%, 0.1% Ga, 0.08% Cu, 0.18% O, 0.06% C, 0.050% N, and the balance Fe. In the accelerated test in this case, the magnet was processed into a size of 8 mm × 8 mm × 2 mm, and the surface thereof was plated with Ni having a thickness of 10 μm, and then the sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity. FIG. 5 shows that the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less is 80% or more and the area of the main phase crystal grains having a crystal grain size of 13 μm or more with respect to the total area of the crystals of the magnet main phase. It can be seen that the corrosion resistance of the R-Fe-B-based sintered permanent magnet having the above specific composition is particularly excellent when the sum of is not more than 10%. Therefore, the size of the crystal grains of the magnet main phase is defined above.
[0012]
The method of controlling the crystal grain size of the main phase of the R-Fe-B sintered permanent magnet having the specific composition to be within the above specified range is not necessarily unique, and various methods or combinations of these methods are not necessarily used. However, in our research, there are considerable difficulties in the ordinary method. Generally, in the production of an R—Fe—B sintered permanent magnet, a raw material coarse powder is pulverized into fine powder, and the fine powder is molded in a magnetic field to obtain a molded product, which is then sintered. A method of forming a sintered body is employed. For example, when fine pulverization is performed using a jet mill, fine powder having a predetermined average particle size and particle size distribution can be obtained by controlling the pressure of gas at the time of pulverization, the supply rate of coarse powder, and the like. If necessary, the particle size distribution of the fine powder can be controlled by performing classification. In forming and sintering the fine powder produced in this way, by selecting an appropriate sintering temperature, time and pattern, the crystal grains of the main phase of the R-Fe-B sintered permanent magnet are selected. It is not always impossible to make the diameter fall within the above specified range. However, it is necessary to set and control many conditions, and it has been found that it is extremely difficult to produce a sintered body having a predetermined crystal grain size with good reproducibility.
[0013]
The present inventors have searched for a method that is easy and suitable for mass production in order to make the crystal grain size of the main phase of the R-Fe-B sintered permanent magnet having the specific composition fall within the above specified range. An R-Fe-B-based quenched ribbon-shaped alloy having a predetermined composition produced by a method called a strip casting method is heat-treated in a predetermined temperature range, and a method of pulverizing the heat-treated alloy into a raw material coarse powder has been found. Further, in crushing the strip-shaped alloy after the heat treatment, it is effective to perform a dehydrogenation treatment after the metal alloy is naturally collapsed by occlusion of hydrogen in order to enhance the fine crushing performance.
FIG. 6 shows Nd 27.8%, Pr 0.45%, Dy 1.7%, B 1.05%, Al 0.05%, Co 2.05%, Ga 0.08%, by weight percentage. It is a figure which shows the cross-sectional structure of the strip-shaped alloy manufactured by the strip cast method which has a composition of 0.09% of Cu, 0.02% of O, 0.004% of N, 0.007% of C, and the balance of Fe ( as cast). A dendrite-like fine structure exists. In the photograph, the phase observed in white corresponds to the main phase of the permanent magnet sintered body having a small amount of rare earth, and the phase observed in black corresponds to the Rrich phase of the permanent magnet sintered body having a large amount of rare earth. Phase. Since this Rrich phase becomes a starting point of destruction at the time of pulverization, when a thin strip alloy in which this Rrich phase is finely dispersed is used as shown in FIG. Easy to generate. Therefore, the sintered body having the magnet main phase particle size distribution of the present invention can be produced relatively easily and with good reproducibility without strictly controlling many conditions during pulverization and sintering. .
However, even if the thin strip alloy (as quenched casting) is directly pulverized as it is to obtain a raw material coarse powder and finely pulverized, it is difficult to obtain a good fine powder particle size distribution. It is difficult to obtain a desired main phase crystal grain size distribution with good reproducibility in the sintered body. The reason for this is that the surface of the ribbon-shaped alloy is hardened by quenching casting, which significantly deteriorates the crushability during fine pulverization.
[0014]
The present inventors have found that, as a means for solving this problem, it is effective to remove the hardening of the surface of the ribbon-shaped alloy by heat-treating the ribbon-shaped alloy in a specific temperature range. The temperature of the heat treatment is set to 800C to 1100C. This is because if the heat treatment temperature is lower than 800 ° C., the removal of the curing is insufficient. Further, at a temperature higher than 1100 ° C., a reaction occurs between the ribbon-shaped alloys during the heat treatment, and it becomes difficult to perform the treatment in a subsequent step. Since the alloy is a thin ribbon containing a large amount of active rare earth elements, the heat treatment needs to be performed in an inert gas atmosphere or in a substantial vacuum. In addition, as described above, after absorbing the hydrogen into the ribbon-shaped alloy after the heat treatment to cause natural collapse, and performing the dehydrogenation treatment, coarsening this is more effective in increasing the pulverizability. is there. This is because, in addition to the effect of removing the hardening of the surface of the ribbon-shaped alloy due to the heat treatment, the effect of embrittlement of mainly the Rrich phase inside the ribbon-shaped alloy due to hydrogen is added.
[0015]
Table 1 shows the main phase crystal of the sintered body when the ribbon-shaped alloy shown in Fig. 6 was heat-treated (1 hr) or pulverized under various conditions to obtain a coarse powder, pulverized under the same conditions, and molded and sintered. This shows the state of the particle size. In addition, the manufacturing method and conditions of the sintered body after the pulverization are described later in detail.
[0016]
[Table 1]

[0017]
Table 1 shows that the sintered body having the main phase particle size distribution of the present invention can be obtained by heat-treating the ribbon-shaped alloy at a temperature of 800 ° C. or higher and using this. Also, as described above, the effectiveness of the hydrogen treatment is clear. At the same time, as shown in Table 1, the state of the main phase grain size in the case where the heat treatment was performed at 700 ° C. was almost the same level as that in the as-quenched state. Thus, it can be seen that at a heat treatment temperature of 700 ° C., the removal of the surface hardened portion of the ribbon-shaped alloy is insufficient.
At the same time, the present inventors have found that by performing a heat treatment on the ribbon-shaped alloy at 800 ° C. or more, Br among the magnetic properties can be particularly improved. The results are also shown in Table 1. From Table 1, Br of the permanent magnet sintered body of the quenched casting state and the ribbon-shaped alloy heat-treated at 700 ° C. is 13.2 to 13.3 KG. On the other hand, when the ribbon-shaped alloy heat-treated at 800 ° C. and 900 ° C. is used, Br rapidly increases to 13.55 KG. At a heat treatment temperature of 1000 ° C., the resulting Br slightly increases to 13.6 KG. At the heat treatment temperatures of 1100 ° C. and 1200 ° C., the increase in Br reaches saturation and remains unchanged at 13.6 KG.
[0018]
7, 8, and 9 show cross-sectional structures of heat treatment temperatures of 900 ° C., 1000 ° C., and 1100 ° C. among the ribbon-shaped alloys shown in Table 1. Comparing these with the quenched as-cast condition (FIG. 6), as the heat treatment temperature increases, both the white structure corresponding to the main phase and the black structure corresponding to the Rrich phase in the ribbon-shaped alloy become coarse. You can see that it is. From these facts, the present inventors have found that a thin strip alloy as-quenched has a fine structure composed of a phase corresponding to the main phase and the Rrich phase. Among the fine powders, those in the polycrystalline state are stochastically present, causing a decrease in the orientation when the fine powders are molded in a magnetic field, resulting in a reduction in Br of the permanent magnet sintered body. Think. At a heat treatment temperature of 700 ° C., the growth of the above structure is insufficient and the orientation cannot be improved. As shown in FIGS. 6 to 9, the internal structure of the ribbon-shaped alloy is coarsened with an increase in the heat treatment temperature. However, this reduces the probability of generation of fine powder in a polycrystalline state, and improves Br. Although it can be considered, as far as judging from the results in Table 1, it is considered that the effect is considerably obtained at the heat treatment temperature of 800 ° C. As the heat treatment temperature of the ribbon-shaped alloy further increases, Br of the obtained sintered body is slightly improved, but tends to be saturated at a heat treatment temperature of 1000 ° C. or more. This is because even when the heat treatment temperature is further increased to promote the coarsening of the structure at the stage where the structure inside the ribbon-shaped alloy has coarsened to some extent and the fine powder in the polycrystalline state has hardly been generated stochastically, It can be understood that it is not reflected as an improvement in Br of the obtained sintered body.
[0019]
As described in detail above, a quenched cast strip alloy having a predetermined composition by a strip casting method is heat-treated in a specific temperature range, or is subjected to hydrogen occlusion to spontaneously disintegrate. By pulverizing, the pulverizability at the time of fine pulverization is improved, and the permanent magnet sintered body manufactured using the same has the main phase crystal grain size of the present invention which is extremely excellent in corrosion resistance. However, it also has high magnetic properties.
The heat treatment temperature of the ribbon-shaped alloy is preferably 800 to 1100 ° C., and the heat treatment time is preferably 15 minutes to 3 hours, more preferably 30 minutes to 3 hours.
[0020]
The reasons for limiting the composition of the R-Fe-B sintered permanent magnet of the present invention will be described below.
The amount of the rare earth element is between 27.0 and 31.0% by weight. The amount of rare earth elements31.0 %If it exceeds 300, the amount of the Rrich phase inside the sintered body increases and the morphology becomes coarse, resulting in poor corrosion resistance. On the other hand, when the amount of the rare earth element is less than 27.0%, the amount of the liquid phase necessary for densification of the sintered body is insufficient, and the density of the sintered body is reduced. iHc decreases together. Therefore, the amount of the rare earth element is set to 27.0 to 31.0%.
The amount of O is between 0.05 and 0.25% by weight. When the amount of O exceeds 0.25%, a part of the rare earth element forms an oxide, and the magnetically effective rare earth element decreases to decrease the coercive force iHc. On the other hand, since the level of the O content of the ingot produced by melting is 0.04% at the maximum, it is difficult to reduce the O content of the final sintered body to this value or less, and the O content is preferably 0.05 to 0.25%. .
The amount of C is between 0.01 and 0.15% by weight. If the amount of C is more than 0.15%, part of the rare earth element forms carbide, and the magnetically effective rare earth element decreases, and the coercive force iHc decreases. The amount of C is more preferably 0.12% or less, further preferably 0.10% or less. On the other hand, the level of the C content of the ingot produced by melting is 0.008% at the maximum, and it is difficult to make the C content of the final sintered body less than this value, and the C content of the sintered body is 0.01 to 0.15%. Is preferred. A specific method for setting the amounts of O and C in the sintered body to the above values will be described later.
[0021]
According to the research results of the present inventors, in order to greatly improve the corrosion resistance of the R—Fe—B based sintered rare earth magnet, it was previously stated that the amount of the rare earth element should be 31.0% or less. It is a necessary condition, but not a sufficient condition, to make the size of the sintered body main phase crystal grain size in the specific range described above. Further, it is necessary to strictly control the amount of N in the sintered body. In the R-Fe-B-based sintered rare earth magnet having the sintered body main phase crystal grain diameter in the specific range and the rare earth amount, O amount, and C amount in the above composition range, the N in the sintered body By setting the (nitrogen) amount in a predetermined range, it is possible to achieve both excellent corrosion resistance and high magnetic properties. The amount of N in the sintered body needs to be 0.02 to 0.15% by weight percentage. The mechanism of the effect of improving the corrosion resistance due to the inclusion of N is not always clear, but N in the sintered body is mainly present in the Rrich phase, and combines with some of the rare earth elements to form a nitride. It is considered that the formation of the nitride suppresses the anodic oxidation of the Rrich phase. When the amount of N is less than 0.02%, the effect of improving the corrosion resistance of the sintered body is not seen, probably because the amount of nitride formed is small. When the amount of N is 0.02% or more, the corrosion resistance of the sintered body improves as the amount of N increases, but when the amount of N exceeds 0.15%, the coercive force iHc sharply decreases. This is considered to be due to the reduction of the magnetically effective rare earth element due to the formation of nitride. For the above reasons, the N amount is set to 0.02 to 0.15%. More preferably, the N content is 0.03 to 0.13%.
[0022]
In the R-Fe-B sintered permanent magnet of the present invention, a part of Fe is Nb, Al, Co, Ga, and Cu.One or more ofThe reason for limiting the replacement amount of each element (in this case, the weight percentage based on the total composition of the permanent magnet after the replacement) will be described below.
The substitution amount of Nb is set to 0.1 to 2.0%. With the addition of Nb, borides of Nb are generated during the sintering process, which suppresses abnormal grain growth. If the Nb substitution amount is less than 0.1%, the effect of suppressing abnormal grain growth of the crystal grains will not be sufficient. On the other hand, if the substitution amount of Nb exceeds 2.0%, the amount of Nb boride generated increases, so that the residual magnetic flux density Br decreases.
The substitution amount of Al is set to 0.02 to 2.0%. The addition of Al has the effect of increasing the coercive force iHc. When the substitution amount of Al is less than 0.02%, the effect of improving the coercive force is small. When the substitution amount exceeds 2.0%, the residual magnetic flux density Br sharply decreases.
The substitution amount of Co is set to 0.3 to 5.0%. The addition of Co improves the Curie point, that is, improves the temperature coefficient of saturation magnetization. When the substitution amount of Co is less than 0.3%, the effect of improving the temperature coefficient is small. When the amount of Co exceeds 5.0%, both the residual magnetic flux density Br and the coercive force iHc sharply decrease.
The substitution amount of Ga is set to 0.01 to 0.5%. The addition of a small amount of Ga improves the coercive force iHc, but when the substitution amount is less than 0.01%, the effect of addition is small. On the other hand, when the substitution amount of Ga exceeds 0.5%, the decrease in the residual magnetic flux density Br becomes remarkable and the coercive force iHc also decreases.
The substitution amount of Cu is set to 0.01 to 1.0%. The addition of a small amount of Cu improves the coercive force iHc, but when the substitution amount exceeds 1.0%, the effect of addition is saturated. When the addition amount is less than 0.01%, the effect of improving the coercive force iHc is small.
[0023]
Next, a method of controlling the N content of the R—Fe—B based sintered permanent magnet, which is a gist of the present invention, will be described.
There are various methods for controlling the N content of the R-Fe-B sintered permanent magnet, and the method is not limited in the present invention and can be selected in the present invention. For example, a raw material powder for R-Fe-B sintered type permanent magnet is charged into a jet mill pulverizer, and then the inside of the jet mill is replaced with Ar gas to substantially reduce the oxygen concentration in the Ar gas. 0%, then N₂A small amount of gas is introduced and N in Ar gas₂Adjust the gas concentration (usually in the range of 0.0001 to 0.1 vol%). This N₂In the process of pulverizing the raw material coarse powder in an Ar gas atmosphere containing a trace amount of gas, mainly the rare earth element in the raw material and N are combined, and the amount of N in the recovered fine powder increases. In collecting the fine powder, a container filled with a solvent such as mineral oil, vegetable oil, or synthetic oil is directly installed in the fine powder collection port of the jet mill, and the fine powder is directly collected in the solvent in an Ar gas atmosphere. The slurry-like raw material thus obtained is wet-formed in a magnetic field to obtain a formed body. The molded body is placed in a vacuum furnace at 5 × 10^-2The molded body is heated to a temperature of about 200 ° C. under a vacuum of about torr to remove the solvent contained in the molded body. Subsequently, the temperature of the vacuum furnace was raised to a sintering temperature of about 1100 ° C.^-4Sintering is performed under a vacuum of about torr to obtain a sintered body. Thus, an R—Fe—B sintered permanent magnet having an O content of 0.25% or less and a C content of 0.15% or less can be obtained. In this case, the control of the amount of N in the sintered body is based on the introduction of N in the Ar gas during the pulverization.₂This is performed by controlling the concentration of gas. The degree of mixing of N into the raw material varies depending on the capacity of the jet mill, the composition and amount of the raw material coarse powder charged, the feed amount of the raw material coarse powder during jet mill pulverization, and the like.
Accordingly, in order to obtain the target amount of N in the sintered body, conditions are determined for each of the conditions at the time of pulverization, and the optimum amount of N in Ar gas is determined.₂It is necessary to determine the gas concentration and pulverize. By such a method, the amount of N in the sintered body can be controlled to 0.02 to 0.15%.
[0024]
The inside of the jet mill is N₂Replace with gas and its N₂The oxygen concentration in the gas is made substantially 0%,₂R-Fe-B system with O content of 0.25% or less, C content of 0.15% or less and N content of 0.02 to 0.15% by finely pulverizing raw material coarse powder in a gas atmosphere A sintered permanent magnet can also be obtained. In this case, the degree of mixing of N into the raw material is controlled by the charged amount of the raw material coarse powder and the feed amount of the raw material coarse powder at the time of pulverization, and a target N amount sintered body is obtained. Since the degree of mixing of N into the raw material varies depending on the type and capacity of the jet mill, conditions are set in advance to set the amount of raw material coarse powder charged and the amount of feed during pulverization. The method of recovering the fine powder after pulverization is in a solvent composed of one or more of mineral oil, vegetable oil, and synthetic oil, and the steps after wet molding are the same as in the case of pulverization in the Ar gas atmosphere. Is the same.
The atmosphere in which the oxygen concentration is substantially 0% as described above means, for example, in the case of a production type mill mill capable of finely pulverizing an R—Fe—B-based raw material coarse powder by about 10 kg / Hr. The atmosphere in which the oxygen concentration in the atmosphere is 0.01 vol% or less, more preferably 0.005 vol% or less, further preferably 0.002 vol% or less in percentage.
By the above method, an R-Fe-B-based sintered body having an O content of 0.25% or less, a C content of 0.15% or less, and an N content of 0.02 to 0.15% can be produced. However, at the same time, by using as a raw material a quenched ribbon-shaped alloy having a predetermined composition which has been subjected to a heat treatment at the temperature range of 800 to 1100 ° C. described above, the main phase crystal grain size in the specific range can be easily obtained. And it can be obtained with good reproducibility.
By subjecting the thus obtained sintered body to heat treatment and working, it is possible to produce an R-Fe-B based sintered permanent magnet having excellent corrosion resistance and high magnetic properties.
[0025]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described specifically with reference to Examples, but the present invention is not limited thereto.
(Example 1)
Nd 27.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.35%, Al 0.08%, Co 2.5%, Ga 0.09% by weight percentage , Cu 0.08%, O 0.03%, C 0.005%, N 0.004%, balance Fe Fe, a strip-shaped alloy having a thickness of 0.2 to 0.5 mm is strip cast. It was produced by the method. This ribbon-shaped alloy was heated at 1000 ° C. for 2 hours in an Ar gas atmosphere. Next, using a hydrogen furnace, the ribbon-shaped alloy was occluded with hydrogen in a hydrogen gas atmosphere at room temperature, and spontaneously collapsed. Next, the ribbon-shaped alloy was heated to 550 ° C. while the inside of the furnace was evacuated, and held at that temperature for 1 hour to perform a dehydrogenation treatment. The collapsed alloy was mechanically crushed in a nitrogen gas atmosphere to obtain a raw material coarse powder of 32 mesh or less.
When the composition of this raw material coarse powder was analyzed, Nd was 27.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.35%, Al 0.08%, Co 2. Analytical values of 5%, 0.09% of Ga, 0.08% of Cu, 0.12% of O, 0.02% of C, 0.008% of N, and the balance Fe were obtained.
After charging 80 kg of the raw material coarse powder into the jet mill, the inside of the jet mill is filled with N.₂Replace with gas, N₂The oxygen concentration in the gas was set to substantially 0% (0.001 vol% as measured by an oxygen analyzer). Then, a pulverization pressure of 7.0 kg / cm²And the raw material coarse powder was pulverized under the conditions of a supply amount of 10 kg / Hr. The average particle size of the fine powder was 3.9 μm.
A container filled with mineral oil (trade name Idemitsu Super Sol PA-30, manufactured by Idemitsu Kosan) is directly installed at the fine powder collection port of the jet mill.₂Fine powder was recovered directly into mineral oil in a gas atmosphere. The raw material after recovery was a raw material slurry in which the fine content of fine powder was 80% by weight by adjusting the amount of mineral oil. 0.8 ton / cm of this raw material slurry is applied to the mold cavity while applying an orientation magnetic field of 12 kOe in the mold cavity.²Wet molding was performed at a molding pressure of The direction of application of the orientation magnetic field is perpendicular to the molding direction. A large number of solvent discharge holes were provided in the upper punch of the mold, and a cloth filter having a thickness of 1 mm was applied to the upper punch surface during molding.
The molded body is 5.0 × 10^-2Heat at 200 ° C. × 1 hour in a torr vacuum to remove mineral oil contained, then 4.0 × 10 4^-4The temperature was increased to 1070 ° C. at a rate of 15 ° C./min under the condition of torr, and the temperature was maintained for 3 hours for sintering.
When the composition of the sintered body was analyzed, Nd 27.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.35%, Al 0.08%, Co 2.5 %, 0.09% of Ga, 0.08% of Cu, 0.16% of O, 0.07% of C, 0.055% of N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less in the total area of the magnet main phase crystals of this sintered body is 93%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is: It was 4%.
This sintered body was subjected to heat treatment once each in an Ar gas atmosphere at 900 ° C. × 2 hours and 480 ° C. × 1 hour. When the magnetic properties were measured after machining, good values as shown in Table 2 were obtained.
In order to evaluate the corrosion resistance of the permanent magnet, the magnet was processed into a fixed size of 8 mm × 8 mm × 2 mm, and its surface was plated with Ni having a thickness of 10 μm. Next, this sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity, and the degree of separation of Ni plating over time was examined. As shown in Table 2, no abnormality was observed in the Ni plating even after lapse of 2500 hours, indicating good corrosion resistance.
[0026]
(Example 2)
Nd 22.3%, Pr 2.0%, Dy 5.5%, B 1.0%, Nb 0.5%, Al 0.2%, Co 2.0%, Ga 0.09% by weight percentage , Cu 0.1%, O 0.02%, C 0.005%, N 0.003%, and the balance of Fe. It was produced by a strip casting method. This ribbon-shaped alloy was heated at 1100 ° C. for 1 hour in an Ar gas atmosphere. Next, using a hydrogen furnace, the thin ribbon-shaped alloy was occluded with hydrogen in a hydrogen gas atmosphere at room temperature and naturally collapsed. Next, the strip-shaped alloy was heated to 550 ° C. while the inside of the furnace was evacuated, and held at that temperature for 1 hour to perform a dehydrogenation treatment. The collapsed alloy was mechanically crushed in a nitrogen gas atmosphere to obtain a raw material coarse powder of 32 mesh or less.
When the composition of this raw material coarse powder was analyzed, 22.3% of Nd, 2.0% of Pr, 5.5% of Dy, 1.0% of B, 0.5% of Nb, 0.2% of Al, and 0.2% of Co. Analysis values of 0%, 0.09% of Ga, 0.1% of Cu, 0.11% of O, 0.02% of C, 0.006% of N, and the balance Fe were obtained.
After 100 kg of the raw material powder is charged into a jet mill, the inside of the jet mill is filled with N.₂Replace with gas, N₂The oxygen concentration in the gas was set to substantially 0% (0.002% as measured by an oxygen analyzer).
. Then, a pulverizing pressure of 8.0 kg / cm²The raw material coarse powder was pulverized under a supply amount of 12 kg / Hr. The average particle size of the fine powder was 3.8 μm.
A container filled with mineral oil (trade name Idemitsu Super Sol PA-30, manufactured by Idemitsu Kosan) is directly installed at the fine powder collection port of the jet mill.₂Fine powder was recovered directly into mineral oil in a gas atmosphere. The raw material after recovery was a raw material slurry having a fine powder content of 77% by weight by adjusting the amount of mineral oil. This raw material slurry was placed in a mold cavity for 1.5 ton / cm while applying an orientation magnetic field of 10 kOe.²Wet molding was performed at a molding pressure of The direction of application of the orientation magnetic field is perpendicular to the molding direction. A large number of solvent discharge holes were provided in the upper punch of the mold, and a cloth filter having a thickness of 1 mm was applied to the upper punch surface during molding.
The molded body is 5.0 × 10^-2Heat at 200 ° C. for 2 hours in a vacuum of torr to remove mineral oil contained, then 5.0 × 10^-4The temperature was raised to 1090 ° C. at a rate of 15 ° C./min under torr conditions, and the temperature was maintained for 3 hours for sintering.
When the composition of the sintered body was analyzed, Nd 22.3%, Pr 2.0%, Dy 5.5%, B 1.0%, Nb 0.5%, Al 0.2%, Co 2.0 %, 0.09% of Ga, 0.1% of Cu, 0.14% of O, 0.06% of C, 0.040% of N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of the total area of the magnet main phase crystals of the sintered body is 95%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is: It was 3%. FIG. 10 shows the relationship between the particle size and the area ratio of this sintered body.
This sintered body was subjected to heat treatment once each in an Ar gas atmosphere at 900 ° C. × 2 hours and 460 ° C. × 1 hour. When the magnetic properties were measured after machining, good values as shown in Table 2 were obtained.
In order to evaluate the corrosion resistance of this permanent magnet, the magnet was processed into a fixed size of 8 mm × 8 mm × 2 mm, and its surface was plated with Ni having a thickness of 10 μm. Next, this sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity, and the degree of separation of Ni plating over time was examined. As shown in Table 2, no abnormality was observed in the Ni plating even after lapse of 2500 hours, indicating good corrosion resistance.
FIG. 10 also shows that the sintered body according to the present invention has a sharp particle size distribution centered on a crystal particle size of 5 μm.
[0027]
(Example 3)
Nd 20.7%, Pr 8.6%, Dy 1.2%, B 1.05%, Al 0.08%, Co 2.0%, Ga 0.09%, Cu 0.1% by weight percentage , O 0.03%, C 0.006%, N 0.004%, and the balance of Fe, a strip-shaped alloy having a thickness of 0.1 to 0.5 mm was produced by a strip casting method. This ribbon-shaped alloy was heated at 900 ° C. for 3 hours in an Ar gas atmosphere. Next, using a hydrogen furnace, the thin ribbon-shaped alloy was occluded with hydrogen in a hydrogen gas atmosphere at room temperature, and naturally collapsed. Next, the ribbon-shaped alloy was heated to 550 ° C. while the inside of the furnace was evacuated, and held at that temperature for 1 hour to perform a dehydrogenation treatment. The collapsed alloy was mechanically crushed in a nitrogen gas atmosphere to obtain a raw material coarse powder of 32 mesh or less.
When the composition of this raw material coarse powder was analyzed, Nd was 20.7%, Pr 8.6%, Dy 1.5%, B 1.05%, Al 0.08%, Co 2.0%, and GaO. The analytical values of 09%, Cu 0.1%, O 0.13%, C 0.03%, N 0.009%, and the balance Fe were obtained.
After 50 kg of the raw material coarse powder was charged into the jet mill, the inside of the jet mill was replaced with Ar gas, and the oxygen concentration in the Ar gas was reduced to substantially 0% (0.002 vol% as measured by an oxygen analyzer). . Next, N in Ar gas₂The gas concentration was set to 0.005 vol%. Then, a pulverizing pressure of 7.5 kg / cm²The raw material coarse powder was pulverized under the condition of a supply amount of 8 kg / Hr.
A container filled with mineral oil (trade name: Idemitsu Super Sol PA-30, manufactured by Idemitsu Kosan Co., Ltd.) was directly installed at the fine powder collecting port of the jet mill, and the fine powder was directly collected into the mineral oil in an Ar gas atmosphere. The raw material after recovery was made into a raw material slurry in which the fine content of fine powder was 75% by weight by adjusting the amount of mineral oil. The average particle size of the fine powder was 4.0 μm. This raw material slurry is placed in a mold cavity while applying an orientation magnetic field of 13 kOe to 0.6 ton / cm.²Wet molding was performed at a molding pressure of The direction of application of the orientation magnetic field is perpendicular to the molding direction. A large number of solvent discharge holes were provided in the upper punch of the mold, and a cloth filter having a thickness of 1 mm was applied to the upper punch surface during molding.
The molded body is 6.0 × 10^-2Heat at 180 ° C. × 4 hours in torr vacuum to remove mineral oil contained, then 3.0 × 10 4^-4The temperature was raised to 1070 ° C. at a rate of 15 ° C./min under torr conditions, and the temperature was maintained for 2 hours for sintering.
When the composition of the sintered body was analyzed, Nd 20.7%, Pr 8.6%, Dy 1.2%, B 1.05%, Al 0.08%, Co 2.0%, Ga 0.09 %, 0.1% Cu, 0.18% O, 0.07% C, 0.075% N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of the total area of the magnet main phase crystals of the sintered body is 90%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is: 5%.
This sintered body was subjected to heat treatment once each in an Ar gas atmosphere at 900 ° C. × 2 hours and 510 ° C. × 1 hour. When the magnetic properties were measured after machining, good values as shown in Table 2 were obtained.
In order to evaluate the corrosion resistance of this permanent magnet, the magnet was processed into a fixed size of 8 mm × 8 mm × 2 mm, and its surface was plated with Ni having a thickness of 10 μm. Next, this sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity, and the degree of separation of Ni plating over time was examined. As shown in Table 2, no abnormality was observed in the Ni plating even after lapse of 2500 hours, indicating good corrosion resistance.
[0028]
(Example 4)
Nd 22.0%, Pr 5.0%, Dy 1.5%, B 1.1%, Al 1.0%, Co 2.5%, O 0.02%, C 0.005% by weight percentage , N 0.005%, and the balance of Fe, a ribbon-shaped alloy having a thickness of 0.1 to 0.4 mm was produced by a strip casting method. This ribbon-shaped alloy was heated at 1000 ° C. for 2 hours in an Ar gas atmosphere. The ribbon-like alloy after the heat treatment was mechanically crushed in a nitrogen gas atmosphere to obtain a raw material coarse powder of 32 mesh or less.
When the composition of this raw material powder was analyzed, Nd was 22.0%, Pr was 5.0%, Dy was 1.5%, B was 1.1%, Al was 1.0%, Co was 2.5%, and O.O. Analysis values of 14%, C 0.01%, N 0.009%, and the balance Fe were obtained. After charging 50 kg of the raw material coarse powder into a jet mill, the inside of the jet mill is N₂Replace with gas, N₂The oxygen concentration in the gas was set to substantially 0% (0.002 vol% as measured by an oxygen analyzer). Then, a pulverization pressure of 7.0 kg / cm²And the raw material coarse powder was pulverized under the conditions of a supply amount of 10 kg / Hr. The average particle size of the fine powder was 4.2 μm.
A container filled with mineral oil (trade name Idemitsu Super Sol PA-30, manufactured by Idemitsu Kosan) is directly installed at the fine powder collection port of the jet mill.₂Fine powder was recovered directly into mineral oil in a gas atmosphere. The raw material after recovery was a raw material slurry in which the fine content of fine powder was 78% by weight by adjusting the amount of mineral oil. This raw material slurry was placed in a mold cavity while applying an orientation magnetic field of 11 kOe to 0.5 ton / cm.²Wet molding was performed at a molding pressure of The direction of application of the orientation magnetic field is perpendicular to the molding direction. A large number of solvent discharge holes were provided in the upper punch of the mold, and a cloth filter having a thickness of 1 mm was applied to the upper punch surface during molding.
The molded body is 5.0 × 10^-2Heat at 200 ° C. × 2 hours in a vacuum of torr to remove the mineral oil contained, then 2.0 × 10^-4The temperature was raised to 1080 ° C. at a rate of 15 ° C./min under torr conditions, and the temperature was maintained for 3 hours for sintering.
When the composition of the sintered body was analyzed, Nd 22.0%, Pr 5.0%, Dy 1.5%, B 1.1%, Al 1.0%, Co 2.5%, O 0.17 %, C 0.07%, N 0.060%, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of the total area of the magnet main phase crystals of the sintered body is 88%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is 7%.
This sintered body was subjected to heat treatment once each in an Ar gas atmosphere at 900 ° C. × 2 hours and 600 ° C. × 1 hour. When the magnetic properties were measured after machining, good values as shown in Table 2 were obtained.
In order to evaluate the corrosion resistance of this permanent magnet, the magnet was processed into a fixed size of 8 mm × 8 mm × 2 mm, and its surface was plated with Ni having a thickness of 10 μm. Next, this sample was left under the conditions of 2 atm, 120 ° C., and 100% humidity, and the degree of separation of Ni plating over time was examined. As shown in Table 2, no abnormalities were observed in the Ni plating even after lapse of 2000 hours, indicating good corrosion resistance.
[0029]
(Comparative Example 1)
The ribbon-shaped alloy produced in Example 1 was directly placed in a hydrogen furnace without performing heat treatment, and was allowed to occlude hydrogen in a hydrogen gas atmosphere at room temperature to be naturally collapsed. Thereafter, dehydrogenation treatment and mechanical crushing were performed under the same conditions as in Example 1 to obtain raw material coarse powder of 32 mesh or less. When the composition of this raw material coarse powder was analyzed, Nd was 27.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.35%, Al 0.08%, by weight percentage. The analysis values were 2.5% Co, 0.09% Ga, 0.08% Cu, 0.10% O, 0.02% C, 0.007% N, and the balance Fe.
This raw material powder was pulverized under the same conditions as in Example 1. The average particle size of the obtained fine powder was 4.4 μm, which was coarser than that of Example 1.
Subsequent steps such as collection of fine powder, preparation of raw material slurry, wet molding, demineralized oil and sintering, heat treatment, and evaluation of corrosion resistance were also performed under the same conditions as in Example 1. When the composition of the sintered body was analyzed, Nd 27.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.35%, Al 0.08%, Co 2.5 %, 0.09% of Ga, 0.08% of Cu, 0.14% of O, 0.06% of C, 0.045% of N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of the total area of the magnet main phase crystals of the sintered body is 78%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is: It was 12%.
When the magnetic properties of this permanent magnet were evaluated, as shown in Table 2, both Br and iHc were slightly lower than the values of Example 1. Further, as shown in Table 2, the corrosion resistance of this permanent magnet was found to be at a level at which no abnormalities were observed in Ni plating even after 1200 hours and no problem occurred in practical use. Slight separation of Ni plating occurred, and it was found that the corrosion resistance was inferior to that of the sintered body manufactured in Example 1.
[0030]
(Comparative Example 2)
An R-Fe-B-based alloy ingot having the same composition as in Example 2 was produced. The composition analysis values of this alloy were 22.3% for Nd, 2.0% for Pr, 5.5% for Dy, 1.0% for B, 0.5% for Nb, 0.2% for Al, 0.2% for Co in terms of percentage by weight. 0%, Ga 0.09%, Cu 0.1%, O 0.01%, C 0.004%, N 0.002%, and the balance Fe. Since precipitation of α-Fe was recognized in the structure of the alloy, the alloy ingot was subjected to a liquefaction treatment at 1100 ° C. for 6 hours in an argon gas atmosphere to eliminate the precipitation. Next, the alloy ingot was placed in a hydrogen furnace, and hydrogen was absorbed at room temperature to cause spontaneous collapse. The alloy after spontaneous collapse was subjected to dehydrogenation treatment and mechanical crushing under the same conditions as in Example 2 to obtain a raw material coarse powder of 32 mesh or less.
When the composition of the raw material coarse powder was analyzed, Nd 22.3%, Pr 2.0%, Dy 5.5%, B 1.0%, Nb 0.5%, Al 0.2%, The analysis values were 2.0% Co, 0.09% Ga, 0.1% Cu, 0.10% O, 0.02% C, 0.005% N, and the balance Fe.
This raw material powder was pulverized under the same conditions as in Example 2. The average particle size of the obtained fine powder was 4.7 μm, which was coarser than that of Example 1.
Subsequent processes such as collection of fine powder, preparation of raw material slurry, wet molding, sintering with demineralized oil, heat treatment, and evaluation of corrosion resistance were also performed under the same conditions as in Example 2. When the composition of the sintered body was analyzed, Nd 22.3%, Pr 2.0%, Dy 5.5%, B 1.0%, Nb 0.5%, Al 0.2%, Co 2.0 %, 0.09% of Ga, 0.1% of Cu, 0.12% of O, 0.06% of C, 0.030% of N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of the total area of the magnet main phase crystals of this sintered body is 61%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is: 22%. FIG. 11 shows the relationship between the particle size and the area ratio of this sintered body.
When the magnetic properties of this permanent magnet were evaluated, as shown in Table 2, the values were as good as those of Example 2. Further, as shown in Table 2, the corrosion resistance of this permanent magnet was found to be at a level at which no abnormalities were observed in Ni plating even after 1000 hours and no problem was found in practical use. Slight separation occurred in a part of the Ni plating, and it was found that the corrosion resistance was inferior to that of the permanent magnet manufactured in Example 2. Further, from FIG. 11, the sintered body of Comparative Example 2 has a broad particle size distribution centered on the crystal grain size of 8 μm, and is 13 μm compared with FIG. 10 which is the particle size distribution of the sintered body of Example 2. It can be seen that the area ratio of the above crystal grain size is large.
[0031]
(Comparative Example 3)
The same raw material powder as used in Example 3 was pulverized under the same conditions as in Example 3. However, N in Ar gas₂No gas was introduced. The average particle size of the fine powder was 4.0 μm. Subsequent steps such as collection of fine powder, preparation of raw material slurry, wet molding, demineralized oil and sintering, heat treatment, and evaluation of corrosion resistance were also performed under the same conditions as in Example 3. When the composition of the sintered body was analyzed, Nd 20.7%, Pr 8.6%, Dy 1.2%, B 1.05%, Al 0.08%, Co 2.0%, Ga 0.09 %, Cu 0.1%, O 0.18%, C 0.07%, N 0.010%, and the balance Fe.
The sum of the areas of the main phase crystal grains having a crystal grain size of 10 μm or less was 92%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more was 4% with respect to the total area of the magnet main phase crystals of this sintered body. .
When the magnetic properties of this permanent magnet were evaluated, good values as shown in Table 1 were obtained. However, it was found that the corrosion resistance of this permanent magnet was not good because Ni plating was separated in 192 hours as shown in Table 2.
[0032]
(Comparative Example 4)
Nd 30.0%, Pr 0.5%, Dy 1.5%, B 1.05%, Nb 0.8%, Al 0.2%, Co 3.0%, Ga 0.08% by weight percentage , Cu 0.1%, O 0.02%, C 0.005%, N 0.005%, balance Fe Fe, a strip-shaped alloy having a thickness of 0.2 to 0.5 mm is strip cast. It was produced by the method. This ribbon-shaped alloy was heated at 950 ° C. for 4 hours in an Ar gas atmosphere. Next, using a hydrogen furnace, the thin ribbon-shaped alloy was occluded with hydrogen in a hydrogen gas atmosphere at room temperature, and naturally collapsed. Next, the ribbon-shaped alloy was heated to 550 ° C. while the inside of the furnace was evacuated, and held at that temperature for 1 hour to perform a dehydrogenation treatment. The collapsed alloy was mechanically crushed in a nitrogen gas atmosphere to obtain a raw material coarse powder of 32 mesh or less.
An analysis of the composition of the raw material coarse powder revealed that Nd was 30.0%, Pr was 0.5%, Dy was 1.5%, B was 1.05%, Nb was 0.8%, Al was 0.2%, and Co was 3.0%. Analysis values of 0%, 0.08% of Ga, 0.1% of Cu, 0.12% of O, 0.02% of C, 0.009% of N, and the balance Fe were obtained.
After 100 kg of the raw material powder is charged into a jet mill, the inside of the jet mill is filled with N.₂Replace with gas, N₂The oxygen concentration in the gas was set to substantially 0% (0.001 vol% as measured by an oxygen analyzer). Then, a pulverizing pressure of 7.5 kg / cm²And the raw material coarse powder was pulverized under the conditions of a supply amount of 10 kg / Hr. The average particle size of the fine powder was 4.1 μm.
A container filled with mineral oil (trade name: Idemitsu Super Sol PA-30, manufactured by Idemitsu Kosan) is directly installed at the fine powder collection port of the jet mill.₂Fine powder was recovered directly into mineral oil in a gas atmosphere. The recovered raw material was made into a raw material slurry having a fine content of 70% by weight by adjusting the amount of mineral oil. 0.8 ton / cm of this raw material slurry was applied in a mold cavity while applying an orientation magnetic field of 14 kOe.²Wet molding was performed at a molding pressure of The direction of application of the orientation magnetic field is perpendicular to the molding direction. A large number of solvent discharge holes were provided in the upper punch of the mold, and a cloth filter having a thickness of 1 mm was applied to the upper punch surface during molding.
The molded body is 5.0 × 10^-2Heat at 200 ° C. × 2 hours in a vacuum of torr to remove mineral oil contained, then 3.0 × 10^-4The temperature was raised to 1080 ° C. at a rate of 15 ° C./min under torr conditions, and the temperature was maintained for 3 hours for sintering.
When the composition of the sintered body was analyzed, Nd was 30.0%, Pr was 0.5%, Dy was 1.5%, B was 1.05%, Nb was 0.8%, Al was 0.2%, and Co was 3.0. %, 0.08% of Ga, 0.1% of Cu, 0.15% of O, 0.07% of C, 0.060% of N, and the balance Fe.
The sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less in the total area of the magnet main phase crystals of the sintered body is 92%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is It was 4%.
This sintered body was subjected to heat treatment once each in an Ar gas atmosphere at 900 ° C. × 2 hours and 550 ° C. × 1 hour. When the magnetic properties were measured after machining, good values as shown in Table 2 were obtained.
In order to evaluate the corrosion resistance of this permanent magnet, the magnet was processed into a fixed size of 8 mm × 8 mm × 2 mm, and its surface was plated with Ni having a thickness of 10 μm. Next, this sample was left under the conditions of 2 atm, 120 ° C. and 100% humidity, and the degree of separation of Ni plating over time was examined. As shown in Table 2, the separation of Ni plating occurred in 48 hours as shown in Table 2. However, it was found to be completely unsuitable for practical use.
[0033]
[Table 2]

[0034]
【The invention's effect】
According to the present invention,Excellent corrosion resistance without deteriorating magnetic properties R-Fe-B A sintered sintered permanent magnet is obtained.
[Brief description of the drawings]
FIG. 1 shows that the sum of the areas of the main phase grains having a crystal grain size of 10 μm or less with respect to the total area of the magnet main phases is 96%, and the sum of the areas of the main phase grains having a crystal grain size of 13 μm or more is 1%. It is a figure which shows the metal structure of a sintered type permanent magnet.
FIG. 2 is a graph showing the sum of the area of main phase crystal grains having a crystal grain size of 10 μm or less to 64% and the total area of main phase crystal grains having a crystal grain size of 13 μm or more to 17% of the total area of the magnet main phase. FIG. 3 is a view showing a metal structure of a tie permanent magnet.
FIG. 3 shows that the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less to the total area of the magnet main phase is 96%, and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is 1%. It is a figure which shows the metallographic structure after 5000 hours of corrosion resistance evaluation tests of a sintered permanent magnet.
FIG. 4 is a graph showing the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less of 64% and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more with respect to the total area of the magnet main phase being 17%. It is a figure which shows the metallographic structure after a lapse of 2000 hours of the corrosion resistance evaluation test of the permanent magnet.
FIG. 5 is a graph showing the ratio of the sum of the areas of the main phase crystal grains having a crystal grain size of 10 μm or less to the total area of the magnet main phase crystals and the crystal grain size of 13 μm or more with respect to the total area of the magnet main phase crystals. FIG. 4 is a diagram showing a relationship between the ratio of the sum of the areas of crystal grains of the main phase and the elapsed time until the start of Ni plating separation in an accelerated corrosion resistance test.
FIG. 6 is a view showing a metal structure of a cross section of a thin strip alloy produced by a strip casting method.
FIG. 7 is a view showing a metal structure of a cross section after heat treatment at 900 ° C. of a thin strip alloy produced by a strip casting method.
FIG. 8 is a view showing a metal structure of a cross section after heat treatment of a ribbon-shaped alloy produced by a strip casting method at 1000 ° C.
FIG. 9 is a view showing a metal structure of a cross section after heat treatment of a ribbon-shaped alloy produced by a strip casting method at 1100 ° C.
FIG. 10 shows that the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less is 95% and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is 3% with respect to the total area of the magnet main phase crystals. FIG. 3 is a view showing the particle size distribution of the sintered body of FIG.
FIG. 11 shows that the sum of the area of the main phase crystal grains having a crystal grain size of 10 μm or less is 61% and the sum of the areas of the main phase crystal grains having a crystal grain size of 13 μm or more is 22% with respect to the total area of the magnet main phase crystals. FIG. 3 is a view showing a particle size distribution of a sintered body as shown in FIG.

Claims

R in weight percentage (R is one or more of the rare earth elements including Y) 27.0-31.0%, B 0.5-2.0%, N 0.02-0.15%, O 0.25% or less, C 0.15% or less, balance A sintered permanent magnet having a composition of Fe,
Manufactured using an R-Fe-B-based quenched alloy for the sintered permanent magnet having a two-phase structure substantially corresponding to the main phase and the Rrich phase of the sintered permanent magnet, thereby improving corrosion resistance. A sintered permanent magnet, characterized in that:

R in weight percentage (R is one or more of rare earth elements including Y) 27.0-31.0%, B 0.5-2.0%, N 0.02-0.15%, O 0.25% or less, C 0.15% or less, balance Fe has a composition, and the sum of the areas of the main phase grains having a crystal grain size of 10 μm or less with respect to the total area of the magnet main phase grains is 80% or more, and the main phase grains having a crystal grain size of 13 μm or more are formed. A sintered permanent magnet having a total area of 10% or less,
Manufactured using an R-Fe-B-based quenched alloy for the sintered permanent magnet having a two-phase structure substantially corresponding to the main phase and the Rrich phase of the sintered permanent magnet, thereby improving corrosion resistance. A sintered permanent magnet, characterized in that:

A part of Fe Nb 0.1~2.0%, Al 0.02~2.0% , Co 0.3~5.0%, Ga 0.01~0.5%, those substituted with one or more of Cu 0.01 to 1.0% according to Item 3. The sintered permanent magnet according to item 1 or 2 .