JP3848906B2 - Method for producing rare earth alloy powder for sintered magnet - Google Patents

Description

【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_２Fe₁₄B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_２Fe₁₄B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_２Fe₁₄B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/cm²、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_２Fe₁₄B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_２Fe₁₄B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_２Fe₁₄B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/cm²、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-yB_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_２T₁₄B相、２ R-rich相、３ B-rich相、４Ａ R_２T₁₄B相内部より派生したR_２T₁₄B粒子、４Ｂ R_２T₁₄B相外縁部より派生したR_２T₁₄B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]
BACKGROUND 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 a raw material for a sintered magnet for an electromagnetic motor used in, for example, home appliances, OA equipment, and industrial machines. The present invention relates to a method for producing a rare earth alloy powder.
[0002]
[Prior art]
A rare earth-transition metal-boron sintered magnet is usually produced by crushing a melt-cast alloy, press-molding it in a magnetic field, and sintering and heat-treating it in a vacuum or an inert gas atmosphere.
FIG. 1 is a schematic diagram showing the internal structure of a rare earth-transition metal-boron casting alloy. In FIG. 1, the internal structure of the cast alloy is mostly composed of R ₂ T ₁₄ B phase 1 with a size of several tens of μm, and non-magnetic R-rich phase 2 and B-rich phase 3 are present at the phase boundary. Distributed. If the amount of these nonmagnetic phases is large, it will cause a decrease in the remanent magnetization of the magnet. Therefore, it is preferable to remove the R-rich phase and the B-rich phase as much as possible. As the removal method, for example, a technique described in JP-A No. 64-48406 is known. According to this, unnecessary R-rich and B-rich phases are removed by subjecting the cast alloy to hydrogen storage / release treatment and making it fine powder, followed by magnetic selection.
[0003]
[Problems to be solved by the invention]
However, according to the method of the same publication, the R-rich phase is attached to the surface of the R ₂ T ₁₄ B particles recovered after the magnetic separation, and the R-rich phase corresponding to the R-rich phase remains. In this regard, the publication accepts that a certain amount of R-rich phase is indispensable for maintaining the coercive force. However, since the R-rich phase is easily combined with oxygen, unnecessary oxygen is taken into the powder in the subsequent magnet manufacturing process. As a result, the incorporated oxygen has an adverse effect on the magnetization as an impurity, and there is a problem that the sinterability is lowered due to oxidation and the residual magnetization is finally lowered. The gazette overlooks this problem.
[0004]
Therefore, the object of the present invention is to reduce the amount of oxygen incorporated into the powder of the cast alloy, avoid the influence of impurities and the decrease in sinterability, and prevent the decrease in remanent magnetization, and the rare earth alloy for sintered magnets. It is in the provision of the manufacturing method of powder.
[0005]
[Means for Solving the Problems]
The invention of claim 1
_{(1) R x T 1-} xy B y alloys (wherein, R Nd, Pr, Dy, 1 or more selected from Tb and Ho, T is Fe, 1 or more selected from 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) pulverizing the heat-treated alloy obtained by the step (2);
(4) A method for producing a rare earth alloy powder for a sintered magnet, comprising a step of selectively collecting particles adsorbed by a permanent magnet or an electromagnet from the powder obtained in the step (3).
The invention of 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 an amorphous or crystal grain of 100 nm or less during the step (3), during the step (4), between the steps (2) and (3), or after the step (4). Constructed microcrystalline RM-TB alloy (wherein RM is one or more selected from rare earth elements that require either Dy or Tb, T is one or more selected from Fe, Co and Ni) 3) The rare earth alloy powder for sintered magnet according to claim 1 or 2, characterized by mixing 5 to 30% by weight with respect to the heat-treated alloy.
[0006]
[Action]
FIG. 2 is a schematic diagram showing the internal structure of the alloy after the (2) heat treatment step 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 subjecting the composition alloy to a heat treatment at 1000 to 1120 ° C. after casting. Reference numeral 2 is an R-rich phase, and 3 is a B-rich phase.
FIG. 3 is a schematic view showing particles derived from 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 In 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 increased due to coarsening, the number of derived particles 4A is 10 times or more that of the particles 4B. Although these particles are mixed after pulverization, when particles adsorbed by a permanent magnet or an electromagnet are taken out, only R ₂ T ₁₄ B particles that are ferromagnetic are selectively recovered. More than 90% of the R ₂ T ₁₄ B powder obtained in this way is R ₂ T ₁₄ B particles 4A, and the R-rich phase hardly adheres to the surface. The volume ratio is greatly reduced compared to the conventional case. As a result, the amount of oxygen taken into the powder is drastically reduced. In the (3) powdering step, it is preferable to powderize to an average particle size of 1 to 15 μm.
Further, in the present invention, by adding a hydrogen storage / release treatment to the cast alloy during the process of pulverizing the heat-treated alloy, embrittlement of the alloy is promoted and the time required for pulverization is shortened. The hydrogen storage / release treatment can be performed, for example, by leaving it in a hydrogen atmosphere at room temperature and pressure for several hours and then heating it at a temperature of 300 to 650 ° C. for several hours under a pressure of 10 ⁻² Torr or less. .
In addition, during the step (3), during the step (4), between the steps (2) and (3), or after the step (4), a microcrystal composed of amorphous or 100 nm or less crystal grains RM-TB alloy (wherein RM is one or more selected from rare earth elements essential for either Dy or Tb, and T is one or more selected from Fe, Co and Ni) By adding a step of mixing 5 to 30% by weight with respect to the heat-treated alloy as a binder, the sintering aid causes the R-rich grain boundary phase in the subsequent sintering and heat treatment steps for magnet production. Therefore, the coercive force is greatly improved by increasing the sinterability and simultaneously diffusing Dy or Tb contained in the sintering aid to the surface of the R ₂ T ₁₄ B crystal. Furthermore, since the sintering aid is amorphous or microcrystalline and is not easily oxidized, unnecessary oxygen is not taken into the powder even when mixed.
[0007]
DETAILED DESCRIPTION OF THE INVENTION
Embodiment 1 FIG.
Dissolving various Nd _x Fe _1-xy B _y alloys having compositions (0.10 ≦ x ≦ 0.35,0.04 ≦ y ≦ 0.17) in the high-frequency vacuum melting furnace, after heating to a temperature of 1300 to 1500 ° C., cast in a water-cooled casting mold Thus, a plate-like cast alloy having a thickness of 10 mm was produced. When the internal structure was observed with an optical microscope, the size of the R ₂ T ₁₄ B phase of each sample was about 30 μm. These alloys were then heated for 8 hours at various temperatures from 900 to 1150 ° C. using a vacuum furnace. As a result, most of the cast alloy reacts violently with the alumina pedestal at temperatures exceeding 1120 ° C, and sound heat treatment is not performed, and at temperatures below 1000 ° C, the R ₂ T ₁₄ B phase is about 50 µm. It grew only, and it turned out that sufficient coarsening does not occur. Therefore, the heat treatment temperature is preferably 1000 to 1120 ° C. FIG. 4 shows the results of examining the internal structure of the alloy when heat treatment was performed in the same temperature range. The x mark in the figure is a composition in which the presence of α-Fe phase or Nd ₂ Fe ₁₇ phase is recognized. These phases are a factor that greatly deteriorates the magnetic properties of the magnet, and are not suitable as raw materials. Composition. Further, the symbol Δ indicates a composition in which the reaction between the alloy and the pedestal in the furnace is so strong that a sound heat treatment cannot be performed, or the R ₂ T ₁₄ B phase is not sufficiently coarsened. On the other hand, a circle indicates a composition in which a heat treatment is performed smoothly and an R ₂ T ₁₄ B phase having a size of 200 μm or more is obtained. It can be said. 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 Six alloy samples of Ni _0.01 B _0.08 were cast in a high frequency vacuum melting furnace. Thereafter, each alloy sample was heated at a temperature of 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 disk mill, and a jet mill, and each of the obtained powder samples was subjected to a permanent magnet type magnetic separator to selectively collect 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. In addition, it is confirmed that the contents of the R-rich phase and B-rich phase in the powder are both less than 0.2% by volume, and the content of the R ₂ Fe ₁₄ B phase exceeds 99.6% by volume. It was done. On the other hand, a Nd _0.38 Fe _0.56 B _0.06 amorphous alloy powder was separately prepared using a liquid quenching apparatus as a sintering aid, and this was added to each sample powder at a rate of 12% by weight and mixed uniformly. Transverse magnetic field molding was performed with a press machine under conditions of a pressure of 600 kgf / cm ² and a magnetic field of 15000 G. Next, sintering was performed in a tube furnace at 1090 ° C in vacuum for 2 hours, followed by rapid cooling, followed by heat treatment in vacuum at 600 ° C for 2 hours to obtain a cylindrical magnet sample having an outer diameter of 10 mm × height of 7 mm. It was. Each sample was magnetized with a 30000 G pulse magnetic field, and a demagnetization curve was measured with a BH curve tracer. Further, the oxygen content was measured by an inert gas melting-infrared absorption method. The measurement results are shown in Table 1. For comparison, the table shows the characteristics of a magnet produced by a conventional method in which heating at 1100 ° C. is not performed after casting and no sintering aid is used. Table 1 shows that the oxygen concentration of the magnet obtained by the present invention is around 1500 ppm, which is less than half that of the magnet of the comparative example. This reflects that the amount of the Nd-rich phase contained in the alloy powder is small in the present invention and the amount of oxidation of the powder is reduced. On the other hand, in the comparative example, since heating at 1100 ° C. is not performed after casting, the crystal grains of the cast alloy are not sufficiently coarsened. For this reason, R ₂ Fe ₁₄ B particles having an R-rich phase attached to the surface during pulverization. There are many derivatives. After the magnetic sorting, the content of the R-rich phase contained in the collected sample powder was examined. As a result, it was found that the content was about 3.0% by volume, much higher than the present invention. It can be said that this large amount of R-rich phase took in a large amount of oxygen during the production process, and as a result, the oxygen content in the comparative example was at a high level close to 5000 ppm. Next, when the internal structure of the magnet after the sintering / heat treatment process 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 when the sintering aid undergoes phase separation during the sintering process, but is oxidized because the sintering aid is amorphous or microcrystalline. There is almost no. Therefore, in the present invention, there is almost no decrease in sinterability due to oxidation of the R-rich phase, and a denser internal structure can be obtained, and at the same time, an amount of R-rich grain boundary necessary for maintaining the holding force can be 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 an effect of preventing a decrease in sinterability. Due to these synergistic effects, as shown in Table 1, the magnet of the present invention has a high residual magnetization compared to the comparative magnet 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-like 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 a temperature of 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 successively pulverized by a jaw crusher, a disk mill, and a jet mill to make a fine powder, and then subjected to an electromagnet magnetic separator to selectively recover 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 prepared using a liquid quenching apparatus as a sintering aid. As a result of observation using an electron microscope, the surface of the ribbon was amorphous, but the inner portion having a thickness of about 15 μm was occupied by crystal grains having a diameter of several tens of nm to 100 nm, and was microcrystalline. The ribbon was sequentially pulverized 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 a transverse magnetic field molding was performed with a magnetic press under a pressure of 600 kgf / cm ² and a magnetic field of 15000 G. Next, each compact was sintered in a tube furnace at 1090 ° C in vacuum for 2 hours, then rapidly cooled, and then heat treated in vacuum at 600 ° C for 2 hours, and a cylindrical magnet having an outer diameter of 10 mm × height of 7 mm. A sample was obtained. Each sample was magnetized with a 30000 G pulse magnetic field, and a demagnetization curve was measured with a BH curve tracer. The measurement results are shown in Table 2. When the weight ratio of the auxiliary powder to the powder sample is less than 5%, the coercive force is less than 7 kOe, and depending on the operating point of the magnet, there is a possibility that a large demagnetization occurs, so that the reliability is lacking. On the other hand, when the weight ratio exceeds 30%, the energy product is significantly reduced. For this reason, the weight ratio of the auxiliary powder to the powder sample is preferably 5 to 30%, 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-like 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 a temperature of 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 pulverized 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 _0.27 Dy _0.11 Fe _0.56 B _0.06 , Nd _0.38 Fe _0.56 B _0.06 , Pr _0.38 Fe using a liquid quenching device as a sintering aid Five alloy ribbons of _0.56 B _0.06 were prepared. The thickness of the ribbon was about 40 μm. When observed using an electron microscope, the inside of the ribbon was amorphous. After mixing each ribbon with a weight ratio of 18% by weight with respect to the coarse powder, each mixed powder was left in a hydrogen atmosphere at room temperature and pressure for 8 hours, followed by 600 ° C. under a pressure of 10 ⁻² Torr. Hydrogen was released by heating at the temperature of 8 hours. Thereafter, the mixture was sequentially pulverized with a disk mill and a jet mill to obtain a fine powder. Although the time spent for jet mill pulverization was about ¼ of the conventional time, the powder was atomized to an average crystal grain size of about 4 μm due to the embrittlement effect by the hydrogen storage and release treatment. Next, each fine powder was subjected to an electromagnet magnetic separator to collect particles. At this time, the fine particles derived from the amorphous alloy ribbon, which is a sintering aid, are also ferromagnetic and thus adsorbed to the electromagnet and recovered by being mixed with Nd ₂ Fe ₁₄ B particles. Each of the obtained mixed powders was subjected to a transverse magnetic field molding with a magnetic field press machine under conditions of a pressure of 600 kgf / cm ² and a magnetic field of 15000 G. Next, each compact was sintered in a tube furnace at 1090 ° C in vacuum for 2 hours, then rapidly cooled, and then heat treated in vacuum at 600 ° C for 2 hours, and a cylindrical magnet having an outer diameter of 10 mm × height of 7 mm. A sample was obtained. Each sample was magnetized with a 30000 G pulse magnetic field, and a demagnetization curve was measured with a BH curve tracer. Table 3 shows the measurement results. Magnets made with a sintering aid containing Dy or Tb gave a larger coercive force than magnets not containing these elements.
[0012]
[Table 3]

[0013]
【The invention's effect】
The invention of claim 1
_{(1) R x T 1-} xy B y alloys (wherein, R Nd, Pr, Dy, 1 or more selected from Tb and Ho, T is Fe, 1 or more selected from 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) pulverizing the heat-treated alloy obtained by the step (2);
(4) A method for producing a rare earth alloy powder for a sintered magnet, characterized by comprising a step of selectively collecting 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 alloy powder, avoid the influence of impurities and the decrease in sinterability, and prevent the decrease in residual magnetization.
[0014]
The invention of 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. Embrittlement is promoted and the time required for pulverization is shortened.
[0015]
The invention according to claim 3 is an amorphous or crystal grain of 100 nm or less during the step (3), during the step (4), between the steps (2) and (3), or after the step (4). Constructed microcrystalline RM-TB alloy (wherein RM is one or more selected from rare earth elements that require either Dy or Tb, T is one or more selected from Fe, Co and Ni) 3) is mixed in a proportion of 5 to 30% by weight with respect to the heat-treated alloy. At the same time, the coercive force improves.
[Brief description of the drawings]
FIG. 1 is a schematic diagram showing the internal structure of a rare earth-transition metal-boron casting alloy.
FIG. 2 is a schematic view showing an internal structure of a rare earth-transition metal-boron casting alloy after heat treatment in the present invention.
FIG. 3 is a schematic diagram showing particles derived from pulverization of a rare earth-transition metal-boron casting alloy after heat treatment in the present invention.
4 is a diagram showing a composition region in which a powder having sound characteristics is obtained in the first embodiment. FIG.
[Explanation of symbols]
1 R ₂ T ₁₄ B phase, 2 R-rich phase, 3 B-rich phase, 4 A R ₂ T ₁₄ B particles derived from inside the R ₂ T ₁₄ B phase, 4 B R ₂ T ₁₄ B phase derived from the outer edge R ₂ T ₁₄ B particles, 5 R-rich particles derived from R-rich phase, 6 B-rich particles derived from B-rich phase.

Claims (3)

_{(1) R x T 1-} xy B y alloys (wherein, R Nd, Pr, Dy, 1 or more selected from Tb and Ho, T is Fe, 1 or more selected from 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) pulverizing the heat-treated alloy obtained by the step (2);
(4) A method for producing a rare earth alloy powder for a sintered magnet, comprising a step of selectively collecting 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 hydrogen storage / release treatment of the heat-treated alloy. During the step (3), during the step (4), between the steps (2) and (3), or after the step (4), it is amorphous or microcrystalline 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 either Dy or Tb is essential, and T is at least one selected from Fe, Co and Ni). The method for producing a rare earth alloy powder for a sintered magnet according to claim 1, wherein the mixture is mixed at a ratio of 5 to 30% by weight.

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