JP6429020B2 - Rare earth permanent magnet - Google Patents

Description

The present invention relates to a rare earth permanent magnet, and more particularly to a rare earth permanent magnet in which a part of R is replaced with a heavy rare earth element in an RTB based sintered magnet.

An R-T-B sintered magnet (R is a rare earth element, T is Fe or Fe partially substituted by Co, and B is boron) having a tetragonal R ₂ T ₁₄ B compound as a main phase is excellent. It is known to have magnetic properties and has been a typical high-performance permanent magnet since the invention in 1982 (Patent Document 1).

An R-T-B type sintered magnet in which the rare earth element R is made of Nd, Pr, Dy, Tb, and Ho has a large anisotropic magnetic field Ha and is preferable as a permanent magnet material. Among these, Nd-Fe-B permanent magnets with rare earth element R as Nd have a good balance of saturation magnetization Is, Curie temperature Tc, and anisotropic magnetic field Ha, and R using other rare earth elements R in terms of resource and corrosion resistance. -It is widely used in consumer, industrial, transportation equipment and the like because it is superior to TB sintered magnets.

At present, it is desired to improve the magnetic characteristics of the RTB-based sintered magnet, and in particular, many ideas have been made to increase the residual magnetic flux density Br and the coercive force HcJ.

JP 59-46008 A

In recent years, rare earth magnets have a wide variety of uses, and higher magnetic properties are required than ever before. In particular, in the application of an R-T-B system sintered magnet to a hybrid vehicle or the like, the magnet is exposed to a relatively high temperature, so it is important to suppress high temperature demagnetization due to heat. In order to suppress this high temperature demagnetization, it is necessary to increase the coercivity of the RTB-based sintered magnet at room temperature.

The present invention has been made in view of such a situation, and an object of the present invention is to provide a permanent magnet capable of giving a coercive force higher than that of conventional magnets to an RTB-based sintered magnet. .

In order to solve the above-described problems and achieve the object, the rare earth-based permanent magnet of the present invention has a main phase of R ₂ T ₁₄ B structure (where R is at least one rare earth element including Y, and T And Fe or B partially substituted by Co, and B is boron) and a sintered body having a grain boundary phase including a Ho—Y—N phase. The cross-sectional area ratio of the Ho—Y—N phase with respect to the grain boundary phase is 2.0% or more and 60% or less, Ho in the Ho—Y—N phase is 5.0 at% or more and 80 at% or less, and Y is 5. It is 0 at% or more and 80 at% or less.

In the present invention, the unit cross section of the sintered body cross section is a 50 μm square region.

The present inventors diligently studied whether there is a method for greatly improving the coercive force in the RTB-based sintered magnet. As a result, it was found that a high coercive force can be obtained by introducing a Ho—Y—N phase as a grain boundary phase. Although the reason is not clear, the present inventors infer as follows. First, the Ho—Y—N phase is stable in energy, hardly dissolves in contact with the main phase or other grain boundary phases, and the grains of the main phase during sintering are only present in the grain boundary phase. It is thought that it has the effect of suppressing growth. Second, the Ho—Y—N phase is energetically stable as a nitride, and its stability is higher than that of the Ho—Y—O phase, so that it is difficult to oxidize, and the amount of oxygen contained in the sintered body is kept low. It is possible. As a result, since oxygen forms an oxide with rare earth elements, it is considered that the concentration of rare earth elements that greatly contribute to magnetism is reduced and the influence of deterioration of magnetic properties can be reduced. Moreover, workability and corrosion resistance are improved because the rare earth oxide can be reduced. When the cross-sectional area ratio of the Ho—Y—N phase in the grain boundary phase is 2.0% or more and 60% or less, the effect of suppressing the grain growth of the main phase particles appears and the coercive force is improved.

Moreover, the composition of the Ho—Y—N phase is such that Ho is 5.0 at% to 80 at% and Y is 5.0 at% to 80 at%, so that rare earth element oxides can be reduced, and magnetic characteristics can be reduced. Can be improved.

As described above, according to the present invention, the RTB-based sintered magnet can have a higher coercive force than before.

Hereinafter, the present invention will be described in detail based on embodiments. In addition, this invention is not limited by the content described in the following embodiment and an Example. In addition, constituent elements in the embodiments and examples described below include those that can be easily assumed by those skilled in the art, those that are substantially the same, and those in a so-called equivalent range. Furthermore, the constituent elements disclosed in the embodiments and examples described below may be appropriately combined or may be appropriately selected and used.

The RTB-based sintered magnet according to the present embodiment contains 11 to 18 at% of rare earth element (R). If the concentration of R is less than 11 at%, the R ₂ T ₁₄ B phase, which is the main phase of the R-T-B sintered magnet, is not sufficiently generated, and α-Fe having soft magnetism is precipitated and retained. The magnetic force is significantly reduced. On the other hand, when R exceeds 18 at%, the volume ratio of the R ₂ T ₁₄ B phase, which is the main phase, decreases, and the residual magnetic flux density decreases.

The RTB-based sintered magnet according to the present embodiment contains 5 to 8 at% of boron (B). When B is less than 5 at%, a high coercive force cannot be obtained. On the other hand, when B exceeds 8 at%, the residual magnetic flux density tends to decrease. Therefore, the upper limit of B is 8 at%.

The RTB-based sintered magnet according to the present embodiment contains 74 to 83 at% of the transition metal element T, and T in the present invention requires Fe, and among these, Co is 4.0 at% or less. Can be contained. Co forms the same phase as Fe, but is effective in improving the Curie temperature and improving the corrosion resistance of the grain boundary phase. Moreover, the RTB-based sintered magnet to which the present invention is applied can contain one or two of Al and Cu in a range of 0.01 to 1.2 at%. By containing one or two of Al and Cu in this range, it is possible to increase the coercive force, increase the corrosion resistance, and improve the temperature characteristics of the obtained sintered magnet.

The RTB-based sintered magnet according to this embodiment allows the inclusion of other elements. For example, elements such as Zr, Ti, Bi, Sn, Ga, Nb, Ta, Si, V, Ag, and Ge can be appropriately contained.

The RTB-based sintered magnet according to this embodiment has R ₂ T ₁₄ B crystal grains (main phase particles), and the size of the main phase particles is about 1 to 10 μm.

The sintered magnet has an R-rich phase (grain boundary phase) in which R is more concentrated than the R ₂ T ₁₄ B crystal grains at a grain boundary formed by two or more adjacent main phase particles. Yes. The R-rich phase includes a Ho—Y—N phase in which R species are Ho and Y and contains nitrogen, a Nd-rich phase in which R species is Nd, a rare earth nitride phase, a rare earth oxide phase, and a boron (B) atom. A B-rich phase having a high concentration may be included.

Hereinafter, preferred examples of the production method of the present invention will be described. In the manufacture of the RTB-based sintered magnet of this embodiment, first, a raw material alloy is prepared so that an RTB-based sintered magnet having a desired composition can be obtained. The raw material alloy can be produced by a strip casting method or other known melting methods in a vacuum or an inert gas, preferably in an Ar atmosphere. In the strip casting method, a molten metal obtained by melting a raw metal in a non-oxidizing atmosphere such as an Ar gas atmosphere is ejected onto the surface of a rotating roll. The melt rapidly cooled by the roll is rapidly solidified in the form of a thin plate or flakes (scales). This rapidly solidified alloy has a homogeneous structure with a crystal grain size of 1 to 50 μm. The raw material alloy can be obtained not only by the strip casting method but also by a melting method such as high frequency induction melting. In order to prevent segregation after dissolution, for example, it can be solidified by pouring into a water-cooled copper plate. An alloy obtained by the reduction diffusion method can also be used as a raw material alloy.

In the case of obtaining an RTB-based sintered magnet in the present invention, the so-called two-alloy method of creating a sintered magnet from two types of alloys, a main phase composition alloy and a grain boundary phase composition alloy, is used as a raw material alloy. Basic.

In the present invention, it is important to precipitate the Ho—Y—N phase at the grain boundary. Therefore, the grain boundary phase composition alloy used as the main phase composition alloy as a two-alloy method must contain Ho and Y.

The raw material alloy is subjected to a grinding process. The main phase composition alloy and the grain boundary phase composition alloy are ground separately or together. The pulverization process includes a coarse pulverization process and a fine pulverization process. First, the raw material alloy is coarsely pulverized until the particle size becomes about several hundred μm. For coarse pulverization, a stamp mill, jaw crusher, brown mill or the like is used. Prior to coarse pulverization, it is effective to perform pulverization by allowing hydrogen to be stored in the raw material alloy and then releasing it. The hydrogen releasing treatment is performed for the purpose of reducing hydrogen as an impurity as a rare earth sintered magnet. The heating and holding temperature for storing hydrogen is 200 ° C. or higher, preferably 350 ° C. or higher. The holding time varies depending on the relationship with the holding temperature, the thickness of the raw material alloy, etc., but is at least 30 minutes or longer, preferably 1 hour or longer. The hydrogen release process is performed with a nitrogen gas flow. Thereby, a Ho-YN phase can be generated.

After the coarse pulverization process, the process proceeds to the fine pulverization process. A jet mill is mainly used for fine pulverization, and a coarsely pulverized powder having a particle size of about several hundreds of μm has an average particle size of 2.5 to 6 μm, preferably 3 to 5 μm. The jet mill releases a high-pressure inert gas from a narrow nozzle to generate a high-speed gas flow, accelerates the coarsely pulverized powder with this high-speed gas flow, collides with the coarsely pulverized powder, and collides with the target or the container wall. It is a method of generating a collision and crushing.

Fatty acids or fatty acid derivatives and hydrocarbons for the purpose of improving lubrication and orientation during molding, such as zinc stearate, calcium stearate, aluminum stearate, stearamide, oleamide, stearic acid or oleic acid Ethylene bisisostearic amide, hydrocarbon paraffin, naphthalene and the like can be added in an amount of about 0.01 to 0.3 wt% during pulverization.

The fine powder is subjected to molding in a magnetic field. What is necessary is just to let the shaping | molding pressure in shaping | molding in a magnetic field be the range of 0.3-3 ton / cm <2> (30-300 Mpa). The molding pressure may be constant from the beginning to the end of molding, may be gradually increased or gradually decreased, or may vary irregularly. The lower the molding pressure is, the better the orientation is. However, if the molding pressure is too low, the strength of the molded body is insufficient and handling problems occur. Therefore, the molding pressure is selected from the above range in consideration of this point. The final relative density of the molded body obtained by molding in a magnetic field is usually 40 to 60%.

The applied magnetic field may be about 10 to 20 kOe (960 to 1600 kA / m). The applied magnetic field is not limited to a static magnetic field, and may be a pulsed magnetic field. A static magnetic field and a pulsed magnetic field can also be used in combination.

Next, the molded body is sintered in a vacuum or an inert gas atmosphere. The sintering temperature needs to be adjusted according to various conditions such as composition, pulverization method, difference in average particle size and particle size distribution, but is sintered at 1000 to 1200 ° C. for 4 to 48 hours. When the sintering time is less than 4 hours, the grain boundary phase composition alloy is melted and cannot be completely spread out, and the main phase particles cannot be wrapped, and particularly the coercive force is adversely affected. . Further, when firing for 48 hours or more, the effect of suppressing grain growth by the Ho—Y—N phase is diminished, the grain growth is remarkably advanced, and particularly the coercive force is adversely affected.

After sintering, the obtained sintered body is subjected to an aging treatment. This process is an important process for controlling the coercive force. Usually, the aging treatment is performed in a vacuum or Ar gas. In the case where the aging treatment is performed in two stages, holding for a predetermined time at around 800 ° C. and around 600 ° C. is effective. When the heat treatment at around 800 ° C. is performed after sintering, the coercive force increases, which is particularly effective in the mixing method. In addition, since the coercive force is greatly increased by the heat treatment at around 600 ° C., the aging treatment at around 600 ° C. is preferably performed when the aging treatment is performed in one stage.

Hereinafter, although the content of the present invention is explained in detail using an example and a comparative example, the present invention is not limited to the following examples. In the present invention, all the production steps are performed in a low oxygen atmosphere in which the oxygen amount is suppressed to 50 ppm or less.

(Example 1)
Main phase composition alloy of 12.2 at% Nd-5.9 at% B-81.9 at% Fe, 6.0 at% Nd-20.0 at% Y-20.0 at% Ho-50.0 at% Fe-2. In order to prepare a grain boundary phase composition alloy of 0 at% Co-1.0 at% Cu-1.0 at% Al, a raw material alloy or thin plate was melted and cast by a strip casting method, each containing a raw material metal or alloy. . Table 1 also summarizes the composition of the grain boundary phase composition alloy. In Table 1, TRE represents the total rare earth concentration of the entire sintered body.

The obtained raw material alloy thin plates were each pulverized with hydrogen to obtain coarsely pulverized powder. Oleic acid amide was added as a lubricant to the coarsely pulverized powder. Subsequently, using an airflow pulverizer (jet mill), fine pulverization was performed in a high-pressure nitrogen gas atmosphere to obtain finely pulverized powder. At this time, nitrogen in the atmosphere is set to 3000 ppm ± 100 ppm.

Subsequently, the main phase composition alloy fine powder and the raised phase composition alloy fine powder were mixed at a ratio of 9: 1 and molded in a magnetic field. Specifically, molding was performed at a pressure of 140 MPa in a magnetic field of 15 kOe to obtain a molded body of 20 mm × 18 mm × 13 mm. The magnetic field direction is a direction perpendicular to the pressing direction. The obtained molded body was fired at 1000 to 1200 ° C. for 4 to 48 hours. Thereafter, an aging treatment is performed at 900 ° C. for 1 hour and at 620 ° C. for 1 hour in an atmosphere in which the nitrogen amount is maintained at 1000 ppm to obtain a sintered body, and an inert gas melting-non-dispersion type infrared absorption method is used. The oxygen content and nitrogen content were measured.

About the obtained sintered compact, the residual magnetic flux density (Br) and the coercive force (HcJ) were measured using the BH tracer. The results were as shown in Table 2.

The obtained sintered body was cut in parallel to the easy axis of magnetization, then embedded in an epoxy resin, and the cross section was polished. For polishing, a commercially available abrasive paper was used, and polishing was performed while changing from a low-grade abrasive paper to a high-grade abrasive paper. Finally, polishing was performed using buffs and diamond abrasive grains. At this time, polishing was performed without adding water or the like. When water is used, the grain boundary phase components are corroded.

After performing the ion milling process on the cross section of the obtained sintered body and removing the influence of the oxide film or nitride film on the outermost surface, the cross section of the R-T-B system sintered magnet is changed to EPMA (Electron Microprobe Analyzer: Electron). Observation and analysis with a Probe Micro Analyzer). Elemental mapping (256 points × 256 points) by EPMA was performed using a 50 μm square region as a unit cross section. The observation position in the cross section is arbitrary. As a result, main phase particles and grain boundaries were distinguished. At least two types of grain boundary phases existed at the grain boundaries. By quantitative analysis, it was confirmed that the two kinds of grain boundary phases were an Nd-rich phase and a Ho-YN phase. Then, the area ratio of the Ho—Y—N phase in the grain boundary phase and the average particle diameter of the main phase particles were determined and shown in Table 2.

A method for calculating the area ratio of the Ho-YN phase in the grain boundary phase will be described below.
(1) The main phase particle part and the grain boundary part were specified from the reflected electron image observed in the unit cross section using an image analysis method, and the area (B) of the grain boundary part was calculated.
(2) The element concentration is calculated from the mapping data of the characteristic X-ray intensities of Ho, Y, and N obtained by EPMA, and each element of Ho, Y, and N in the main phase particle portion specified in (1) above is calculated. The average concentration and standard deviation were calculated.
(3) At the grain boundary, the element concentration is calculated from the mapping data of the characteristic X-ray intensity obtained by EPMA, and the (average value + 3 × standard deviation) of the element concentration in the main phase particle portion obtained in the above (2) The portion where the element concentration value was larger than the value was specified for Ho, Y, and N. And the specified part was defined as a part in which the concentrations of Ho, Y, and N were distributed deeper than the main phase particles.
(4) The part where the grain boundary specified in (1) above and the part in which the concentration of each element of Ho, Y, N specified in (3) above is more densely distributed than in the main phase particle overlaps, It specified as the Ho-YN phase in a grain boundary, and calculated the area (A) of the part.
(5) By dividing the area (A) of the Ho—Y—N phase calculated in (4) above by the area (B) of the grain boundary calculated in (1) above, Ho—Y— occupying the grain boundary. The area ratio (A / B) of the N phase was calculated.

The main phase particle part and the grain boundary part were specified from the backscattered electron image, and the areas of all main phase particles confirmed within a 50 μm square field were calculated one by one. The calculated area was converted into an equivalent circle diameter to obtain main phase particle diameters of all main phase particles. This series of operations was performed with 20 fields of view, and the average value of all the main phase particle diameters obtained was defined as the average particle diameter.

(Example 2)
The composition of the grain boundary phase alloy is 12.0 at% Nd-17.0 at% Y-17.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

Example 3
The composition of the grain boundary phase alloy is 20.0 at% Nd-13.0 at% Y-13.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 1)
Prepared in the same manner as in Example 1 except that the composition of the grain boundary phase alloy is 46.0 at% Nd-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. And evaluated.

(Comparative Example 2)
The composition of the grain boundary phase alloy is 2.0 at% Nd-22.0 at% Y-22.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 3)
Except that the composition of the grain boundary phase alloy is 4.0 at% Nd 21.0 at% Y-21.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al Were prepared and evaluated in the same manner as in Example 1.

(Comparative Example 4)
The composition of the grain boundary phase alloy is 22.0 at% Nd-12.0 at% Y-12.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 5)
The composition of the grain boundary phase alloy is 24.0 at% Nd-11.0 at% Y-11.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

表１は使用した粒界相合金の組成を示している。比較例１〜５及び実施例１〜３を見ると、用いる粒界相合金の組成におけるＮｄの元素比率を下げることにより、Ｈｏ−Ｙ−Ｎ相面積率を上昇させることができた。そしてＨｏ−Ｙ−Ｎ相の面積率の上昇に伴い、酸素含有量は減少し、窒素含有量は上昇した。窒素含有量が増加していることはＨｏ−Ｙ−Ｎ相の焼結体を占める比率が上昇していることに起因する。酸素量の低下は、先述の通りＨｏ−Ｙ−Ｎ相がエネルギー的に安定で、酸化しにくい材料であることに起因すると考えられる。なお、比較例１〜５及び実施例１〜３について、生成したＨｏ−Ｙ−Ｎ相におけるＨｏ、Ｙの元素比率は共に４１ａｔ％であった。Ｈｏ−Ｙ−Ｎ相の面積率が２．０％以上６０％以下の実施例１〜３では、高い保磁力を得ることができた。これはＨｏ−Ｙ−Ｎ相による粒成長抑制効果、及び含有酸素量の低下に起因するものと考えられる。Ｈｏ−Ｙ−Ｎ相の面積率が６２％以上の比較例２〜３では、Ｂｒ、ＨｃＪ共に低下している。これはＨｏ、Ｙの濃度が高すぎために、主相にＨｏ，Ｙが取り込まれたことによるものと考えられる。Ｈｏは主相に入ると、異方性磁界を下げることは無いが飽和磁化を大きく低下させ、Ｙは主相に入ると、異方性磁界と飽和磁化の両方を大きく低下させる。一方でＨｏ−Ｙ−Ｎ相の面積率が１．０％以下の比較例４〜５では、保磁力が十分に上昇していない。これはＨｏ−Ｙ−Ｎ相のもつ粒成長抑制効果を発現するには面積率が不十分だったことが原因と考えられる。

Table 1 shows the composition of the grain boundary phase alloy used. In Comparative Examples 1 to 5 and Examples 1 to 3, the Ho—Y—N phase area ratio could be increased by lowering the elemental ratio of Nd in the composition of the grain boundary phase alloy used. And with the increase in the area ratio of the Ho-Y-N phase, the oxygen content decreased and the nitrogen content increased. The increase in the nitrogen content is attributed to an increase in the proportion of the Ho—Y—N phase sintered body. The decrease in the oxygen amount is considered to be caused by the fact that the Ho—Y—N phase is energetically stable and hardly oxidized as described above. In Comparative Examples 1 to 5 and Examples 1 to 3, the element ratios of Ho and Y in the generated Ho—Y—N phase were both 41 at%. In Examples 1 to 3 in which the area ratio of the Ho—Y—N phase was 2.0% to 60%, a high coercive force could be obtained. This is considered to be caused by the grain growth suppressing effect by the Ho—Y—N phase and the decrease in the oxygen content. In Comparative Examples 2 to 3 in which the area ratio of the Ho—Y—N phase is 62% or more, both Br and HcJ are lowered. This is considered to be because Ho and Y were taken into the main phase because the Ho and Y concentrations were too high. When entering the main phase, Ho does not lower the anisotropic magnetic field, but greatly reduces the saturation magnetization. When entering the main phase, Y greatly reduces both the anisotropic magnetic field and the saturation magnetization. On the other hand, in Comparative Examples 4 to 5 in which the area ratio of the Ho—Y—N phase is 1.0% or less, the coercive force is not sufficiently increased. This is presumably because the area ratio was insufficient to exhibit the grain growth inhibitory effect of the Ho-YN phase.

(Comparative Example 6)
The composition of the grain boundary phase alloy is 12.0 at% Nd-2.0 at% Y-32.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 7)
The composition of the grain boundary phase alloy is 12.0 at% Nd-4.0 at% Y-30.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Example 4)
The composition of the grain boundary phase alloy is 12.0 at% Nd-6.0 at% Y-28.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Example 5)
The composition of the grain boundary phase alloy is 12.0 at% Nd-17.0 at% Y-17.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Example 6)
The composition of the grain boundary phase alloy is 12.0 at% Nd-28.0 at% Y-6.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 8)
The composition of the grain boundary phase alloy is 12.0 at% Nd-30.0 at% Y-4.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

(Comparative Example 9)
The composition of the grain boundary phase alloy is 12.0 at% Nd-32.0 at% Y-2.0 at% Ho-50.0 at% Fe-2.0 at% Co-1.0 at% Cu-1.0 at% Al. Except for this, it was produced and evaluated in the same manner as in Example 1.

比較例６〜９及び実施例４〜６で、粒界相合金におけるＨｏとＹの濃度を変えることにより、Ｈｏ−Ｙ−Ｎ相におけるＨｏとＹの濃度が変化した。そして粒界相全体でＨｏ−Ｙ−Ｎ相が占める面積率はいずれも３０％であった。また、Ｈｏ−Ｙ−Ｎ相におけるＹの比率が上がると、酸素含有量が増加した。これは、Ｙ−Ｎ相はＨｏ−Ｙ−Ｎ相に比べて酸化しやすく、潤滑剤由来の炭素とも結びついてＹ−ＣＯＮ相を形成しやすいことが原因として考えられる。Ｈｏが十分な濃度であることにより、Ｈｏ−Ｙ−Ｎ相の安定性が保たれる。実施例４〜６を見ると、Ｈｏ、Ｙの濃度はともに５〜８０ａｔ％であり、高いＢｒ、ＨｃＪを得られている。これはＨｏ−Ｙ−Ｎ相による粒成長抑制効果、酸素量低減の効果によるものと考えられる。比較例６〜７を見ると、ＢｒとＨｃＪ共に低い値である。これは酸素含有量増加が原因と考えられる。比較例８〜９もまたＢｒとＨｃＪが低い値となっている。Ｈｏ−Ｙ−Ｎ相においてＹの濃度が減少すると、Ｙの代わりにＮｄを取り込みやすくなる。そのために磁性に寄与するＮｄがＨｏ−Ｙ−Ｎ相に奪われ、特性が悪化することが原因と考えられる。

In Comparative Examples 6 to 9 and Examples 4 to 6, the Ho and Y concentrations in the Ho-YN phase were changed by changing the Ho and Y concentrations in the grain boundary phase alloy. The area ratio occupied by the Ho-YN phase in the entire grain boundary phase was 30%. Moreover, the oxygen content increased as the Y ratio in the Ho—Y—N phase increased. This is presumably because the YN phase is more easily oxidized than the Ho-YN phase, and is easily combined with the lubricant-derived carbon to form the Y-CON phase. When the Ho concentration is sufficient, the stability of the Ho-YN phase is maintained. Looking at Examples 4 to 6, the concentrations of Ho and Y are both 5 to 80 at%, and high Br and HcJ are obtained. This is considered to be due to the effect of suppressing grain growth and the effect of reducing the amount of oxygen by the Ho—Y—N phase. Looking at Comparative Examples 6 to 7, both Br and HcJ are low values. This is thought to be due to an increase in oxygen content. In Comparative Examples 8 to 9, Br and HcJ are low values. If the concentration of Y decreases in the Ho-YN phase, it becomes easier to incorporate Nd instead of Y. Therefore, it is considered that Nd contributing to magnetism is lost to the Ho—Y—N phase and the characteristics deteriorate.

(Comparative Example 10)
It was produced and evaluated in the same manner as in Example 5 except that the pulverization step was performed in an Ar gas atmosphere in a low nitrogen atmosphere with nitrogen of 50 ppm or less.

比較例１０ではＨｏ−Ｙ−Ｎ相が確認できなかった。Ｈｏ−Ｙ−Ｎ相の存在により主相粒子の粒成長を抑えられた実施例２と比べると、比較例１０の粒子径は大きくなっており、主相粒子が粒成長し、それに伴い保磁力は低下している。以上から、窒素雰囲気による微粉砕及び時効熱処理工程を経ることによりＨｏ−Ｙ−Ｎ相を生成し、粒成長抑制効果が発現し、磁気特性が向上した。

In Comparative Example 10, the Ho—Y—N phase could not be confirmed. Compared with Example 2 in which the grain growth of the main phase particles was suppressed due to the presence of the Ho—Y—N phase, the particle diameter of Comparative Example 10 was larger, and the main phase particles grew, and the coercive force accordingly. Is falling. From the above, the Ho—Y—N phase was generated through the fine pulverization in the nitrogen atmosphere and the aging heat treatment step, the effect of suppressing grain growth was exhibited, and the magnetic properties were improved.

As described above, the RTB-based sintered magnet according to the present invention has a high residual magnetic flux density and a high coercive force, and is used for consumer, industrial, transportation equipment, etc. that require high output and high efficiency. It is suitable as a permanent magnet.

Claims (1)

The main phase of the R ₂ T ₁₄ B structure (where R is at least one rare earth element including Y, T is Fe or Fe partially substituted by Co, and B is boron), Ho-Y A sintered body having a grain boundary phase including -N phase, and the cross-sectional area ratio of the Ho-YN phase to the total grain boundary phase per unit sectional area in the sintered body is 2.0% or more and 60% or less. A rare earth-based permanent magnet, wherein Ho in the Ho-Y-N phase is 5.0 at% to 80 at% and Y is 5.0 at% to 80 at%.