JP2020155762A - Rare earth magnet and manufacturing method thereof - Google Patents

Abstract

To provide a rare earth magnet having main phase grains of ThMn12 type compound with improved coercive force while suppressing the decrease in magnetization.SOLUTION: The rare earth magnet has main phase crystal grains including Sm and Y of ThMn12 compound and the grain boundary phase including Cu. The main phase crystal grains have a ThMn12 type core phase and a ThMn12 type shell phase. The Sm concentration in the shell phase is higher than the Sm concentration in the core phase.SELECTED DRAWING: Figure 1A

Description

The present invention relates to a rare earth magnet having main phase crystal grains of a ThMn type ₁₂ compound containing yttrium (Y) and a method for producing the same.

In recent years, there has been a demand for the development of magnets with a reduced content of rare earth elements. As used herein, the rare earth element means at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanoids. Here, lanthanoid is a general term for 15 elements from lanthanum to lutetium. As a ferromagnetic alloy in which the composition ratio of rare earth elements contained is relatively small, RT ₁₂ (R is at least one rare earth element, T is Fe, Co or Ni) having a body-centered tetragonal ThMn ₁₂ type crystal structure is used. Are known. Although RT ₁₂ has a high magnetization, it has a problem that the crystal structure is thermally unstable.

Patent Document 1 discloses a rare earth permanent magnet in which a part of Fe, which is a T element, is partially replaced by Ti, which is a structural stabilizing element, to improve thermal stability in exchange for high magnetization. There is.

Patent Document 2 states that by partially substituting the R element of an RFe _12- based compound with an element such as Zr or Hf, the amount of the structural stabilizing element Ti or the like that replaces the transition metal element is reduced to achieve saturation magnetization. A rare earth permanent magnet in which the ThMn ₁₂ structure is stabilized while being maintained is disclosed.

Further, Patent Document 3 discloses an R'-Fe-Co-based ferromagnetic alloy in which Y or Gd is selected as a part of the R element of RFe ₁₂ , and this R'-Fe-Co-based ferromagnetic alloy is disclosed. It is described that the alloy has a ThMn _12- type crystal structure produced by an ultra-quenching method and thus exhibits high magnetic properties.

Further, Patent Document 4 discloses a rare earth permanent magnet in which the stability of the ThMn ₁₂ structure is improved by adding Ti or the like to Patent Document 3.

Further, Patent Document 5 describes that the addition of Cu produces a phase having a non-magnetic and low melting point of 1-4 composition, which enables sintering and high coercive force.

Further, in Patent Document 6, the Sm ₅ Fe ₁₇ system phase, the SmCo ₅ system phase, the Sm ₂ O ₃ system phase, and the Sm ₇ Cu ₃ system, which are different from the ThMn ₁₂ type as the sub-phase with respect to the main phase ThMn ₁₂ , are described. It is described that a high coercive force can be achieved by including at least one of the phases. Further, Non-Patent Document 1 describes that a stable ThMn _12- type crystal structure can be obtained by structural stabilizing elements (Ti, V, Nb, Ta, Mo, W).

Japanese Unexamined Patent Publication No. 64-76703 Japanese Unexamined Patent Publication No. 4-322406 Japanese Unexamined Patent Publication No. 2015-156436 JP-A-2018-125512 Japanese Unexamined Patent Publication No. 2001-189206 JP-A-2017-112300

K. H. J. Buschou, Rep. Prog. Phys. 54 (1991) 1123-1213.

The rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to element substitution of Fe by Ti, but since the amount of Fe substitution by Ti is large, the magnetization is reduced by that amount, and sufficient magnetic properties are provided. I can't get it.

On the other hand, in the rare earth permanent magnet described in Patent Document 2, although the ThMn ₁₂ structure is stabilized by substituting the transition metal element with Ti or the like, the soft magnetic bcc (Fe, Co, Ti) phase And Laves phase are easily generated, and sufficient coercive force cannot be obtained.

The R'-Fe-Co-based ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with a structure-stabilizing element, so that high magnetization, large magnetic anisotropy, and high Curie temperature can be obtained. Since it is a non-equilibrium phase, the main phase compound may decompose in a densification process at high temperature such as sintering.

Unlike Patent Document 2, the rare earth magnet described in Patent Document 4 can suppress the formation of a soft magnetic phase, but it cannot be said that the coercive force is sufficient.

In the rare earth magnet described in Patent Document 5, the magnetic property value may not be high because the amount of Ti added is large.

In the rare earth magnet described in Patent Document 6, when the rare earth rich subphase Sm ₇ Cu ₃ is used, there is a concern that the equilibrium state shifts to a rare earth rich composition than the main phase during the heat treatment and the main phase ratio decreases. Will be done. Further, the coercive force of the sub-phase determines the coercive force of the entire bulk, and it cannot be said that the main phase ThMn ₁₂ itself contributes to the coercive force.

The embodiment of the present disclosure provides a rare earth magnet having a main phase crystal grain of a ThMn ₁₂ type compound, which has an improved coercive force while suppressing a decrease in magnetization.

In an exemplary embodiment, the rare earth magnet of the present disclosure is a rare earth magnet having a main phase crystal grain of an Fe-based ThMn ₁₂ type compound containing Sm and Y and a grain boundary phase containing Cu, and the main component thereof. The phase crystal grains have a ThMn ₁₂ type core phase and a ThMn ₁₂ type shell phase, and the Sm concentration in the shell phase is higher than the Sm concentration in the core phase.

In certain embodiments, the main phase crystal grains contain at least one M element selected from the group consisting of Ti, Si, Al, and Ga, and the concentration of the M element in the shell phase is the core phase. Higher than the concentration of M element in.

In certain embodiments, Si is contained as the M element, and at least the gradient of Si concentration within the core phase is represented by a position-dependent linear function.

Method for producing a rare earth magnet of the present disclosure, in an exemplary embodiment, the composition formula _{R1 1-x R2 x (Fe} 1-y Co y) in w _Ti z Cu _α, R1 has at least Y, further Gd R2 has at least Sm and is at least one selected from the group consisting of La, Ce, Nd and Pr, and x, y, z and w are 0. Bulk of alloy for rare earth magnets, 5 ≦ x ≦ 1.0, 0 ≦ y ≦ 0.4, 8 ≦ w ≦ 12, 0.42 ≦ z <0.70, 0.40 ≦ α ≦ 0.70 A step of preparing a body, a step of preparing a diffusion alloy having a composition formula R1 _1-χ R2 _χ Cu _ω- based diffusion alloy in which χ> x, 2 ≦ ω ≦ 4, and the bulk body and the diffusion. By heating in contact with the alloy and diffusing the component elements of the diffusion alloy into the bulk body, the main phase crystal of the ThMn ₁₂ type compound in which the Sm concentration in the shell phase is higher than the Sm concentration in the core phase. Includes a step of forming a rare earth magnet with grains.

In certain embodiments, the diffusion alloy is represented by the composition formula R1 _1-χ R2 _χ (Cu _1-υ-τ X _υ sis _τ Zn _σ ) _ω , where X is at least one of Al or Ga and χ. > X, 0 <υ≤0.3, 0 <τ <0.5, 0 <σ <0.2, 2≤ω≤4.

According to the embodiment of the present invention, it is possible to provide a rare earth magnet having a main phase crystal grain of a ThMn ₁₂ type compound, which has an improved coercive force while suppressing a decrease in magnetization.

It is a figure which shows the structure in Experimental Example 9 of this disclosure. It is a figure which shows the structure in Experimental Example 11 of this disclosure. It is a figure which shows the line analysis result of the composition in Experimental Example 9 of this disclosure. It is a figure which shows the line analysis result of the composition in Experimental Example 11 of this disclosure. It is a figure which shows the nuclear density distribution of the 3D (100) plane, (110) plane derived by MEM. It is a figure which shows the correspondence relation of each site at the time of the structural change from TbCu ₇ type to ThMn ₁₂ type seen in a Y-Fe compound.

<Rare earth magnet>
The rare earth magnet of the present disclosure is a rare earth magnet having a main phase crystal grain of an Fe-based ThMn ₁₂ type compound containing Sm and Y and a grain boundary phase containing Cu. Further, the main phase crystal grains have a ThMn ₁₂ type core phase and a ThMn ₁₂ type shell phase, and the Sm concentration in the shell phase is higher than the Sm concentration in the core phase. That is, in the present disclosure, the shell phase refers to a phase having a higher Sm concentration than the central portion (core phase) of the main phase crystal grains. Further, in the present disclosure, the whole of the plurality of main phase crystal grains is simply referred to as "main phase", and the phase existing at the grain boundary of the main phase crystal grains is referred to as "grain boundary phase" as a whole.

In an exemplary embodiment, the main phase crystal grains of the rare earth magnet of the present disclosure contain at least one M element selected from the group consisting of Ti, Si, Al, and Ga, and the M element in the shell phase. The concentration of is higher than the concentration of M element in the core phase.

Sm is expected to increase the magnetic anisotropy energy, and M elements such as Ti and Si are expected to have the effect of increasing the magnetic anisotropy magnetic field and reducing stress and strain due to the decrease in magnetization. Entering the main phase of these elements contributes to an increase in coercive force. On the other hand, the introduction of Sm and M elements means that the amount of Y decreases and the structural stability is impaired. Therefore, it is necessary to increase the amount of M elements to compensate for the structural instability. Significantly reduces grain magnetization. That is, the introduction of Sm and M element creates a trade-off relationship between magnetization and coercive force. Therefore, by forming a shell phase in which Sm or both Sm and M elements are concentrated on the outer periphery of the main phase crystal grain, which is the starting point of magnetization reversal, the coercive force is suppressed while suppressing the decrease in magnetization of the magnet as a whole. Improvement is possible.

The main phase crystal grains are preferably in a form in which the shell phase thinly covers the core phase. In particular, when magnetization reversal is likely to occur from a specific crystal plane, it is desirable to form a thin shell phase only on that crystal plane. By forming the thickness of these shell phases at a thickness equal to or greater than the region that is the starting point of magnetization reversal, the effect of sufficiently improving the coercive force is brought about. The scale at which stress and strain are relaxed at the grain ends differs depending on the crystal system, but in general, 1 μm is sufficient, and even a thickness of several nm can exert a sufficient effect.

The Sm concentration in the shell phase is higher than the Sm concentration in the core phase. Such a configuration is realized, for example, by diffusing a metal element from the outside into a bulk body of an alloy for a rare earth magnet. Naturally, it is realized even for samples that are not sufficiently densified, such as magnetic powder and temporary molded bodies. However, in order to obtain a rare earth magnet constructed from the main phase crystal grains of the thin shell phase, it is important to suppress the diffusion of elements into the core phase and efficiently form only the shell phase. For that purpose, it is sufficient if dissolution and reprecipitation occur at a temperature above which the grain boundary phase has sufficient fluidity, and if the temperature is too high, element diffusion into the core phase and enlargement of the shell phase occur. It becomes difficult to realize. Therefore, by preparing a bulk body such as a molded body or a sintered body in which the alloy for rare earth magnets is densified and a diffusion alloy, and diffusing the diffusion alloy along the grain boundary with the bulk body of the alloy for rare earth magnets. It is preferable to make a rare earth magnet. This will be described below. On the other hand, there is also a method of constructing a Sm-rich shell by utilizing the fact that Y contained in the main phase is very easily oxidized. In the heat treatment process for producing a bulk body such as a densified molded body or a sintered body, Y is deprived from the main phase by oxidation at the outer peripheral portion of the magnetic powder, so that a Sm-rich shell is produced. In this case, the main phase needs to have a Y-rich composition in advance.

<Method of manufacturing rare earth magnets>
Before explaining the method for producing a rare earth magnet, the composition of the rare earth magnet alloy and the diffusion alloy will be described.

[Composition of alloy for rare earth magnets]
Rare earth alloy magnet of the present disclosure, in an exemplary embodiment, non-limiting, represented by a composition formula of RT _w Ti _z Cu _α.

In formula _RT w Ti _z Cu _α of the rare-earth magnet alloy, R is R is scandium (Sc), yttrium (Y), and is at least one selected from the group consisting of lanthanoids, T is Fe or Fe And Co. The ranges of the composition ratios w, z, and α are 8 ≦ w ≦ 12, 0.42 ≦ z <0.70, and 0.40 ≦ α ≦ 0.70, respectively. Hereinafter, this alloy may be referred to as "R-Fe-Co-Ti-Cu based rare earth magnet alloy". Further, a structure in which the powder of this alloy is densified may be referred to as a "bulk body".

More specifically, the alloy for R-Fe-Co-Ti- Cu -based rare earth magnet of the present embodiment is described by formula _{R1 1-x R2 x (Fe} 1-y Co y) w Ti z Cu α .. R1 has at least Y and may further have Gd. R2 has at least Sm and may further contain at least one selected from the group consisting of La, Ce, Nd and Pr excluding Gd. x, y, z, w are 0.5 ≦ x ≦ 1.0, 0 ≦ y ≦ 0.4, 8 ≦ w ≦ 12, 0.42 ≦ z <0.70, 0.40 ≦ α, respectively. The composition ratio satisfies ≦ 0.70.

The alloy in such an embodiment contains Sm as R2, but further contains Sm by intergranular diffusion from the outside. Sm can improve the magnetic anisotropy, which is important for increasing the coercive magnetic force.

By adding Cu, a grain boundary phase that coexists with the main phase and is richer in rare earths than the main phase is generated. By forming this rare earth-rich grain boundary phase, the following effects effective for obtaining anisotropic magnetic powder can be obtained.
-The heat treatment facilitates the progress of crystal growth of the main phase.
-It is possible to reduce the formation of different phases during dissolution and solidification.
-Since the rare earth-rich grain boundary phase absorbs and releases hydrogen, cracks occur between the main phase and the grain boundary phase, and the particles can be efficiently pulverized into single crystal particles.

The composition ratios of R1, R2 and Ti determine the magnetic properties of the main phase and high temperature stability. From the viewpoint of magnetic anisotropy, it is desirable that the composition ratio (x) of R2 is equal to or higher than the composition ratio (1-x) of R1. Therefore, it is desirable that 0.5 ≦ x ≦ 1.0. Further, it is desirable that Ti is as small as possible from the viewpoint of saturation magnetization, but it is desirable that Ti be as large as possible from the viewpoint of high temperature stability. As a result of studies by the present inventors, a range of 0.42 ≦ z <0.70 is desirable. In addition, 50 mol% or less of Ti may be replaced with V, W, Nb, Mo and the like.

From the viewpoint of improving the magnetization and magnetic anisotropy at the practical temperature accompanying the increase in the magnetic moment and the increase in the Curie temperature, it is preferable to replace a part of Fe with Co. However, if the amount of substitution is too large, it is not desirable because it causes a decrease in magnetization and magnetic anisotropy. Specifically, the Co substitution amount y is preferably 0 ≦ y ≦ 0.4, and more preferably 0.1 ≦ y ≦ 0.3. When y is about 0.2, the magnetic characteristics are maximized. Since the Co concentration in the shell phase tends to increase slightly due to grain boundary diffusion, the Co function can be efficiently performed in the shell phase by designing the composition of the rare earth magnet alloy of the present invention to y <0.2. It can be expressed.

The amount of Cu determines the amount of grain boundary phase formed in the rare earth magnet alloy. If the amount of grain boundary phase is small, bcc (Fe, Co, Ti) and Nd ₃ (Fe, Ti) ₂₉ type crystal structure phase (hereinafter, 3-29 phase) and Th ₂ Zn ₁₇ at the time of dissolution and solidification Not only can the different phases such as the mold, Th ₂ Ni ₁₇ mold, and the phase of the crystal structure having a partially irregular substitution site inside (hereinafter, 2-17 phase) not disappear, but also anisotropic sintered magnetic powder can be obtained. It is not easy to grow crystals to a size sufficient for. Grain boundary diffusion inside the magnet body is not easy. Further, when the amount of the grain boundary phase is large, the ratio of the main phase decreases, so that the magnetization as a magnet body decreases. As a result of examination by the present inventors, a range of 0.40 ≦ α ≦ 0.70 is desirable.

The grain boundary phase produced is a Cu group. The grain boundary phase is primarily the phase of the crystal structure of KHG ₂ type (hereinafter, 1-2 phase), and other R and Cu, Fe, the ratio of Co 1: Composition of 4 (hereinafter, 1-4 Composition) phase is also included. The molar ratio of R2 / (R1 + R2) of the R element constituting the grain boundary phase is higher than that of the main phase in both the 1-2 phase and the 1-4 composition phase. In addition, Fe and Co are slightly dissolved, and the amount of the solid solution is larger in the phase having a 1-4 composition than in the phase 1-2. Ti hardly dissolves in both phases.

The appropriate amount of w varies depending on the amount of Cu added, but is 8 ≦ w ≦ 12. This is because if w is too large, a soft magnetic bcc (Fe, Co, Ti) phase is generated, and if w is too small, a 2-17 phase or a 3-29 phase is generated. In particular, when constructing a Sm-rich shell by Y oxidation, w is preferably a relatively low value, and 8 ≦ w ≦ 10 is appropriate.

The R-Fe-Co-Ti-Cu based rare earth magnet alloy thus obtained contains a main phase having a ThMn _12- type crystal structure. The TnMn type ₁₂ compound phase in the alloy according to the embodiment of the present invention can typically be stably present even at 1000 ° C. or higher, and is suitably used for adopting a high-performance magnet manufacturing process such as a sintering method. be able to. Generally, the "ThMn _12- type crystal structure" is a tetragonal crystal, but in the embodiment of the present invention, the crystal lattice of the tetragonal crystal is slightly distorted to have orthorhombic symmetry, or in the crystal. Even if the periodicity of the atom is slightly disturbed, it is regarded as "ThMn ₁₂ type crystal structure". Further, when the Ti element concentration is less than 7.7 at% regardless of the presence or absence of Cu, a rare earth element and two Fe elements (Fe dumbbells) may be irregularly substituted inside the crystal. Specifically, the axial direction of the Fe dumbbell is parallel to the c-axis, and the ratio of irregular replacement is 5% at the maximum. This also applies to structural stabilizing elements other than Ti element, and ThMn _12- type crystals stabilized with a smaller amount of stabilizing element than the substitution composition range of stabilizing element described in Non-Patent Document 1 by using Y. It is commonly observed in structures.

Zn or Sn may be further added to the rare earth magnet alloy. It brings about lowering the melting point of the grain boundary phase and forming a uniform two-particle grain boundary, and it becomes possible to sufficiently supply the diffusion alloy to the inside of the bulk body by the grain boundary diffusion. There is also an effect that can be produced by reducing the amount of Cu solid solution in the main phase grains.

[Composition of diffusion alloy]
As a result of diligent research by the present inventors, it is preferable that the composition of the diffusion alloy according to the embodiment of the present invention is Sm richer than the composition of the grain boundary phase of the alloy for rare earth magnets. By introducing a diffusion alloy along the melted grain boundary phase, a 1-12 phase equilibrated with the Sm-rich grain boundary phase is constructed on the outer periphery of the main phase by melting and reprecipitation. Therefore, when the diffusion alloy composition is expressed as R1 _1-χ R2 _χ Cu _ω- based diffusion alloy, it is desirable that χ> x. ω is 2 ≦ ω ≦ 4.

Further, by adding at least Al or Ga to the above diffusion alloy, it is possible to construct a Sm, Ti-rich shell phase in the deep part of the bulk body. More specifically, Al and Ga diffuse along the grain boundary phase, but since they replace a part of the 3d element of the main phase and diffuse into the grain, Fe, which is a constituent element of the main phase, is corresponding to that amount. , Co, Ti and Cu are discharged to the grain boundary phase. In particular, since Al and Ga are selectively coordinated to the 8i site of the main phase (see Non-Patent Document 1), Ti having a strong 8i site orientation remains in the main phase and is concentrated, and Fe is discharged. That is, the amount of Al or Ga correlates with the amount of Fe element discharged into the grain boundary phase, and it is appropriate that the Fe element is an amount that can be absorbed in the grain boundary phase. Further, when the amount of Al or Ga is increased, the thickness of the produced shell phase tends to increase. From the above, when the diffusion source composition is expressed by R1 _1-χ R2 _χ (Cu _1-υ X _υ ) _ω (X = Al, Ga), the range of 0 <υ≤0.3 is appropriate. More preferably, 0.1 ≦ υ ≦ 0.2. Further, the more desirable composition of ω may differ between the addition of Al and the addition of Ga depending on the presence or absence of Si or Zn. In the case of Al, 2.8 <ω ≦ 4 is more desirable, and in the case of Ga, 2 ≦ ω. <2.8 may be more desirable.

Further, further enriching the M element of the shell phase not only further enhances the thermal stability of ThMn ₁₂ type, but also improves the coercive force because the magnetic anisotropy magnetic field is further improved. is important. According to Non-Patent Document 1, V can replace 4/12 of the T element and Si can replace 2/12 of the T element at the maximum. However, when the grain boundary of V is narrow, it is easy to form a compound with the Fe element, and it is difficult to introduce a sufficient amount of V into the deep part of the bulk body by grain boundary diffusion. This also applies to Ti, Cr, Nb, Mo, W and the like. On the other hand, it was found that Si can introduce a sufficient amount into the deep part of the bulk body along the grain boundaries. Si can be suitably used to construct a shell in which the M element is more concentrated in the deep part of the bulk body. If the amount of Si introduced is too large, the amount of R-Cu-Si compound produced becomes too large, and sufficient diffusion into the deep part of the bulk body is hindered. R1 _1-χ R2 _χ (Cu _1-υ-τ X _υ Si _τ ) When the composition of the diffusion alloy is expressed in _ω , the range of 0 <τ <0.5 is appropriate. Further, Si is characterized by having a higher affinity for Cu and a smaller atomic diameter than Fe, and an element satisfying at least a part of these characteristics has a grain boundary in a bulk body having a grain boundary phase containing Cu. A diffusion process can be preferably used.

Therefore, by combining the above three, the main phase crystal grains have a ThMn12 type core phase and a ThMn12 type shell phase, and a shell phase richer in Sm, Ti, and Si than the core phase is constructed. .. In that case, the diffusion source composition is R1 _1-χ R2 _χ (Cu _1-υ-τ X _υ Si _τ ) _ω (X = Al, Ga, χ> x, 0 <υ ≦ 0.3, 0 <τ < 0.5, 2 ≦ ω ≦ 4) is desirable. Further, adding Zn to the diffusion source composition improves the wettability between the main phase and the grain boundary phase, and makes it possible to introduce more metal elements contained in the diffusion source to the deep part of the bulk body. desirable. Zn tends to be unevenly distributed in the grain boundary phase, and hardly diffuses into the main phase. The diffusion source composition in this case is R1 _1-χ R2 _χ (Cu _1-υ−τ X _υ Si _τ Zn _σ ) _ω (X = Al, Ga, χ> x, 0 <υ ≦ 0.3, 0 < The range of τ <0.5, 0 <σ <0.2, 2 ≦ ω ≦ 4) is desirable.

<Step A> Step of preparing a bulk body of an alloy for rare earth magnets A bulk body of an alloy for rare earth magnets (R-Fe-Co-Ti-Cu based alloy for rare earth magnets) having the above-mentioned composition is prepared. As a method for producing an alloy for R-Fe-Co-Ti-Cu based rare earth magnets, known methods such as a mold casting method, a centrifugal casting method, a strip casting method, and a liquid ultra-quenching method can be adopted. In order to suppress the formation of bcc (Fe, Co, Ti) phases, which are particularly unfavorable as raw material alloys for magnets, during solidification of the molten alloy, the strip casting method or liquid ultra-quenching method, which has a relatively high cooling rate, is used. It is appropriate to adopt. When the cooling rate at the time of solidification is slow, the grain size of the precipitated different phase becomes large, so that the different phase is unlikely to disappear in the next heat treatment step. Even nanocrystals generated by the liquid ultra-quenching method can be easily grown into crystal grains of 10 μm or more suitable for obtaining anisotropic magnetic powder by undergoing the next heat treatment step, so that the raw material alloy can be cooled as quickly as possible. It is better to solidify at a rate. It is better that the coagulated structure has less macrosegregation and is as uniform as possible. For example, when produced by the liquid ultra-quenching method, the roll peripheral speed is preferably 1 to 40 m / s in an inert atmosphere.

By applying heat treatment to the alloy for rare earth magnets in the embodiment of the present invention, it is possible to reduce the heterogeneous phase generated in the solidification process and crush the powder composed of single crystal-like particles useful as a raw material for anisotropic sintered magnets. The crystal grains that can be easily obtained by the method can be coarsened. Although it depends on the composition, the eutectic temperature of the 1-2 phase and the 1-4 composition phase is around 820 ° C, the melting point of the 1-2 phase is around 860 ° C, and the melting point of the 1-4 phase is around 880 ° C. is there. Therefore, the heat treatment temperature is preferably 900 ° C. or higher and 1250 ° C. or lower, and more preferably 1000 ° C. or higher and 1100 ° C. or lower. The heat treatment time depends on the temperature, but is preferably 5 minutes or more and 50 hours or less. If the time is too short, a sufficient reaction will not occur to eliminate the heterogeneous phase or the grain growth will be insufficient, and if the time is too long, the rare earth elements will evaporate and oxidize, and the operational efficiency will be poor. In this temperature range, the grain boundary phase becomes a liquid phase and a part of the main phase is dissolved and reprecipitated. Therefore, in the main phase, crystal grains grow dramatically and solidify due to a rapid chemical reaction via the liquid phase. The different phase of time can be easily eliminated if the grain size is not large. The amount of grain boundary phase is preferably 3 wt% or more and 10 wt% or less. If the amount of grain boundary phase is small, not only the heterogeneous phase during dissolution and solidification cannot be eliminated, but also it is not easy to grow the crystal to a size sufficient to obtain anisotropic sintered magnetic powder, and the amount of grain boundary phase. This is because the ratio of the main phase decreases when the amount is large, so that the magnetization as a magnet body decreases. The bcc (Fe, Co, Ti) phase that could not be completely eliminated in this step can be eliminated in step B if the amount is small.

The grain boundary phase contained in this alloy causes hydrogen absorption and release phenomena reflecting the rare earth-rich composition, and is particularly remarkable when heat-treated in hydrogen. For example, hydrogen absorption occurs at temperatures between 250 ° C. and 400 ° C. and hydrogen is released between 540 ° C. and 660 ° C. Therefore, the alloy for rare earth magnets is heated to 400 ° C. or higher in hydrogen to absorb hydrogen, and then the atmosphere is switched to a vacuum atmosphere to sufficiently release hydrogen. In that case, the temperature for switching to the vacuum atmosphere is 700 ° C. or lower. By absorbing and releasing hydrogen, the rare earth rich phase undergoes volume expansion and contraction, and cracks occur between the main phase crystal grains and the subphase. This increases the probability that the magnetic powder will crack at the cracks during the crushing process using a jet mill, stamp mill, ball mill, etc., making it possible to obtain highly orientable anisotropic magnetic powder containing a large amount of fine powder in single crystal units. Become.

The anisotropic magnetic powder thus produced is molded in a magnetic field and heat-treated at 900 ° C. to 1200 ° C. for 5 minutes or more and 50 hours to obtain a densified bulk body. At that time, in order to suppress the evaporation of Sm, a method of simultaneously adding Sm-based alloy fine powder to the ampoule is known and may be adopted. Further, a metal or alloy fine powder that is more easily oxidized because Y is easily oxidized may be added at the same time as an oxygen getter. In order to construct a Sm-rich shell by positively utilizing the fact that Y is easily oxidized, the composition of the entire alloy may be designed on the Y-rich side. In addition, a densified bulk body may be obtained by hot molding or the like. The embodiment of the present invention is preferably used in a densified bulk body composed of a main phase and a grain boundary phase, and the effect thereof may or may not be the orientation of the magnetic powder. Although the effect of improving the coercive force can be found even in the magnetic powder that is not densified, it is not suitable because it is incompatible with the volume magnetization as a bulk body. In that sense, it is more appropriate for the bulk body to be oriented. Further, it may contain a small amount of oxide phase, a phase derived from unavoidable impurities, and the like.

If there is a foreign substance on the surface of the densified bulk body, it may hinder the intergranular diffusion in step C. Therefore, by grinding the surface of the bulk body before the grain boundary diffusion heat treatment, a sufficient coercive force improving effect can be obtained.

<Step B> Step of preparing a diffusion alloy A diffusion alloy having the above-mentioned composition is prepared. The diffusion alloy may be amorphous or composed of a plurality of phases, but it is preferably not extremely segregated. Extreme segregation may hinder diffusion into the bulk body. Therefore, it is appropriate to adopt the strip casting method, the liquid ultra-quenching method, the gas atomizing method, the disc atomizing method, or the arc melting method, which have a relatively high cooling rate. Further, a homogeneous alloy may be prepared by a ball mill, a planetary ball mill, or the like. In addition, a nanoparticle synthesis method using a chemical liquid phase method such as a chemical reaction precipitation method, an inverse micelle method, a hydrothermal synthesis method, or a sol-gel method may be used.

<Step C> A step of forming a rare earth magnet having a main phase crystal grain of a ThMn ₁₂ type compound in which the Sm concentration in the shell phase is higher than that in the core phase. The diffusion alloy prepared in the step B was prepared in the step A. Heat treatment is performed for introduction into the bulk body. Since the diffusion alloy has good wettability with the main phase, if the bulk body has a thickness of about 1 mm, by arranging a sufficient amount in contact with the bulk body at least, the bulk body is placed along the grain boundary phase. It can be sufficiently diffused inside. If the bulk body has a thickness of more than 1 mm, a diffusion source is laid on the lower surface of the bulk body. Furthermore, if it is intended to increase the coercive force at the end of the bulk body, the place where the component elements of the diffusion alloy are diffused will have a high coercive force. Processing etc. is required and may be performed as appropriate.

It is appropriate that the diffusion alloy is arranged between 1 wt% and 80 wt% with respect to the bulk body weight. If the amount of diffusion alloy is small, the effect of improving coercive force cannot be sufficiently obtained, and if the amount of diffusion alloy is large, the shape of the bulk body is deformed and the amount of excess alloy that does not enter the bulk body increases, which is economical. There is no good tendency, such as not being targeted.

The heat treatment temperature is preferably 500 ° C. or higher and 1000 ° C. or lower. More preferably, it is 700 ° C. or higher and 1000 ° C. or lower. If the temperature is too low, the diffusion alloy and the grain boundary phase will not dissolve and a sufficient amount of diffusion alloy cannot be introduced inside the bulk body, and if the temperature is too high, the diffusion of the bulk body into the crystal grains will be remarkable. Therefore, the decrease in magnetization of the entire bulk becomes remarkable. Although it depends on the thickness of the bulk body, the heat treatment time can be set from 5 minutes to 50 hours. If it is too short, sufficient diffusion into the deep part of the bulk body does not occur, and if it is too long, the effect of improving the coercive force tends to be saturated and the operational efficiency is poor.

Further, if the heat treatment is performed again at a temperature lower than the diffusion heat treatment temperature, the coercive force may be improved and it may be effective.

<Step D> Step of Grinding the Surface of the Bulk Body The diffused bulk body may be ground by a known processing method. There is a residue of the diffusion alloy on the surface of the bulk body that has undergone the diffusion heat treatment, and there may be many unfavorable phases such as 2-17 phase and 3-29 phase at the grain boundaries near the diffusion alloy in terms of magnetic characteristics. By grinding the outer peripheral portion of the bulk body by about 10 μm to 500 μm, such a low magnetic characteristic region can be removed, and the magnetic characteristics of the entire bulk can be improved. Since the grinding powder can be reused by applying appropriate treatment, it does not significantly impair economic efficiency. A plating step or the like may be added after this depending on the intended use, and the present invention is not limited by the subsequent steps.

Hereinafter, experimental examples of the present invention will be specifically described, but the present invention is not limited to these experimental examples.

[Experimental example]
<Step A> (Step of preparing a bulk body of an alloy for rare earth magnets)
The raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed step by step in consideration of evaporation of rare earth elements at the time of dissolution. After being sufficiently dissolved in the alumina crucible, the molten metal was discharged on a roll made of Cu rotating at a peripheral speed of 15 m / s. The prepared ultra-quenching thin band was included in the Nb foil, and heat treatment was performed at 1050 ° C. for 1 hour in Ar flow. It was pulverized with a jet mill to obtain magnetic powder having an average particle size of 10 μm or less. The pressed body molded in a 1T magnetic field was put into a furnace and heat-treated at 1150 ° C. for 4 hours. A bulk body of an alloy for rare earth magnets having a thickness of 4 mm was produced by cutting and polishing the surface of the bulk body. The material structure was evaluated using SEM. The composition of the main phase was Y _0.37 Sm _0.63 (Fe _0.83 Co _0.17 ) _11.08 Al _0.07 Si _0.02 Ti _0.57 Cu _0.26 . Al and Si are derived from the crucible used in this step.

<Step B> (Step to prepare diffusion alloy)
The diffusion alloy was produced using an ultra-quenching device (NEV-A30023 manufactured by Nissin Giken Co., Ltd.). A raw material with a purity of 99.9% or higher, which is weighed in consideration of evaporation of rare earth elements, is sufficiently dissolved in a quartz hot water pipe, and then the hot water is discharged on a Cu roll rotating at 20 m / s to form a ribbon. Diffusion alloy was prepared. Specifically, a diffusion alloy having the composition shown in Table 1 below was prepared.

<Step C> (Step of forming a rare earth magnet having main phase crystal grains of ThMn ₁₂ type compound in which the Sm concentration in the shell phase is higher than the Sm concentration in the core phase)
In a container made of SUS, the lower layer of the diffusion alloy, the bulk body of the rare earth magnet alloy, and the upper layer of the diffusion alloy were arranged in this order and then covered with Nb foil. The weights of the lower and upper diffusion alloys were 10% of the weight of the bulk body, respectively. That is, a 20 wt% diffusion alloy was used with respect to the bulk body. The rare earth magnets of Experimental Examples 1 to 13 were prepared by heat-treating at 900 ° C. for 4 hours in Ar flow.

A portion of the sample obtained in step C from the surface to a depth of 100 μm was removed with a surface grinder (step D). The material structure was also evaluated using SEM. In addition, after magnetizing the bulk body and the rare earth magnets of Experimental Examples 1 to 13 with a pulsed magnetic field, the magnetic characteristics at room temperature were evaluated using VSM. In Table 1, the difference in magnetization with respect to the bulk body (magnetization of the experimental example-magnetization of the bulk body) is ΔMr (%), and the difference in coercive force with respect to the bulk body (coercive force of the experimental example-coercive force of the bulk body) is ΔHc. Shown in (%). As shown in Table 1, in all the experimental examples, the coercive force could be improved while suppressing the decrease in magnetization.

For composition analysis, Experimental Examples 1 to 5 and Experimental Examples 9 to 13 were subjected to a desktop scanning electron microscope (JCM manufactured by JEOL Ltd.) to which an energy dispersive X-ray spectrometer (EX-37001 manufactured by JEOL Ltd.) was attached. -6000 Plus), Experimental Examples 6 to 8 are electric field emission scanning electron microscopes (JSM-7001F manufactured by JEOL Ltd.) to which an energy dispersive X-ray spectrometer (JED-2300F manufactured by JEOL Ltd.) is attached. Was used respectively. Table 1 shows the composition near the center of the shell phase in the main phase crystal grains of Experimental Examples 1 to 13. The composition of the main phase crystal grains of Examples 1 to 13 near the center of the core phase is Y _0.37 Sm _0.63 (Fe _0), which is the composition of the main phase of the bulk body of the rare earth magnet alloy. _.83 Co _0.17 ) _11.08 Al _0.07 Si _0.02 Ti _0.57 Cu _0.26 was equivalent.

The magnetic characteristics were evaluated at room temperature with an electromagnet type vibrating sample magnetometer (VSM, VSM-5-20 manufactured by Toei Kogyo Co., Ltd.) after performing pulse magnetization of 6T. The absolute value of magnetization was corrected with Ni.

In Experimental Example 1, the Sm-Cu grain boundary phase was formed by diffusion. It was confirmed that the main phase crystal grains had a ThMn ₁₂ type core phase and a ThMn ₁₂ type shell phase, and a shell phase richer in Sm than the core phase was formed. However, since the formation of the shell could not be confirmed in the deep part of the bulk body, the coercive force was increased, but _{ΔH cJ} was smaller than that in the other experimental examples. This is due to the formation of the Sm rich shell and the disappearance of a small amount of the bcc (Fe, Co, Ti) phase that could not be completely removed in step A.

In Experimental Example 2, an Al-containing Sm—Cu-based grain boundary phase was formed by diffusion. It was confirmed that a Sm-rich shell phase was generated even in the deep part of the bulk body. The average composition of the shell phase was Y _0.18 Sm _0.82 (Fe _0.83 Co _0.17 ) _10.58 Al _0.51 Si _0.02 Ti _0.62 Cu _0.27 . The main phase crystal grains had a ThMn ₁₂ type core phase and a ThMn ₁₂ type shell phase, and it was confirmed that a shell phase having a composition Sm richer and Ti slightly richer than the core phase was constructed. Since the composition was richer in rare earth elements than the equilibrium composition, a 3-29 phase shell was also produced. Therefore, preferably, 2.8 <ω ≦ 4, which is a rare earth element poorer than the actual composition, is suitable. In fact, in Experimental Example 3, when ω = 3.4 was prepared, the formation of the 3-29 phase was not observed.

In Experimental Example 4, the Sm—Cu-based grain boundary phase contained in Ga was formed by diffusion. A shell phase having an average composition of Y _0.20 Sm _0.80 (Fe _0.82 Co _0.18 ) _10.60 Al _0.26 Si _0.13 Ti _0.71 Cu _0.30 was generated in the deep part of the bulk body. .. The main phase crystal grains had a ThMn _12- type core phase and a ThMn _12- type shell phase, and it was confirmed that a shell phase having a composition richer in Sm, Si, Ga, and Ti than the core phase was constructed. On the other hand, since bcc (Fe, Co, Ti) was generated at the grain boundaries, the coercive force was lowered. This is because the composition was poorer than the equilibrium composition. Therefore, 2 ≦ ω <2.8, which is rich in rare earth elements, is more suitable than the actual composition. Actually, in Experimental Example 5, when ω = 2.3 was prepared, no formation of bcc (Fe, Co, Ti) was observed. In addition, Al, which is an unavoidable impurity mixed in step A, is concentrated in the shell phase.

In Experimental Example 6, a Sm—Cu-based grain boundary phase containing Al and Si was formed by diffusion. Sm (Si _0.5 Cu _0.5 ) ₂ was newly produced in the diffusion alloy with a main composition. If the amount of the R-Cu-Si compound produced is too large, sufficient diffusion into the deep part of the bulk body is hindered. Si dissolves 5 to 10 at% in the Cu portion in the 1-2 phase and hardly dissolves in the 1-4 phase. Since Sm (Si _0.5 Cu _0.5 ) ₂ is produced, the alloy composition of the 1-2 phase and the 1-4 phase is on the rare earth element poor side. Therefore, it was substantially 2.8 <ω ≦ 4 shown in Experimental Example 2, and the 3-29 phase was not confirmed. The average composition of the shell phase produced in the deep part of the bulk body is Y _0.20 Sm _0.80 (Fe _0.82 Co _0.18 ) _10.60 Al _0.26 Si _0.13 Ti _0.71 Cu _0.30 It became. The main phase crystal grains had a ThMn _12- type core phase and a ThMn _12- type shell phase, and it was confirmed that a shell phase having a composition richer in Sm, Si, and Ti than the core phase was constructed. The shell phase was clearly observed even in the deeper part of the bulk body than in Experimental Examples 2 and 3.

In Experimental Example 7, a Sm—Cu-based grain boundary phase containing Ga and Si was formed by diffusion. The constituent phase in the diffusion alloy is the same as in Experimental Example 6, although there is a difference between Ga and Al. Therefore, it was substantially 2.8 <ω ≦ 4, and bcc (Fe, Co, Ti) was generated at the grain boundaries. Therefore, in Experimental Example 8, when ω = 2.3 was prepared, no formation of bcc (Fe, Co, Ti) was observed at the grain boundaries. The average composition of the shell phase was Y _0.20 Sm _0.80 (Fe _0.82 Co _0.18 ) _10.75 Al _0.09 Si _0.19 Ti _0.61 Cu _0.31 Ga _0.05 . .. It was confirmed that a shell phase having a composition richer in Sm and Si than in the core phase and slightly rich in Ga and Ti was constructed. The shell phase was clearly observed even in the deeper part of the bulk body than in Experimental Examples 4 and 5.

In Experimental Example 9, a Sm—Cu-based grain boundary phase containing Al, Si, and Zn was formed by diffusion. The average composition of the shell phase was Y _0.20 Sm _0.80 (Fe _0.82 Co _0.18 ) _10.73 Al _0.26 Si _0.10 Ti _0.64 Cu _0.25 Zn _0.02 . .. However, Zn falls into the category of identification error. It was confirmed that a shell phase having a composition richer in Sm, Si, Al, and Ti than the core phase was constructed. The shell phase was clearly observed even in the deeper part of the bulk body than in Experimental Example 6.

In Experimental Example 10, a Sm—Cu-based grain boundary phase containing Ga, Si, and Zn was formed by diffusion. The constituent phases in the diffusion alloy are the same as in Experimental Example 7, although there is a difference in whether or not they contain Zn. It was substantially 2.8 <ω ≦ 4, and bcc (Fe, Co, Ti) was generated at the grain boundaries. Therefore, in Experimental Example 11, when ω = 2.3 was prepared, no formation of bcc (Fe, Co, Ti) was observed. The average composition of the shell phase was Y _0.20 Sm _0.80 (Fe _0.82 Co _0.18 ) _10.82 Al _0.03 Si _0.10 Ti _0.71 Cu _0.25 Ga _0.09 . .. It was confirmed that a shell phase having a composition richer in Sm, Si, Ga, and Ti than the core phase was constructed. The shell phase was clearly observed even in the deeper part of the bulk body than in Experimental Examples 7 and 8.

In Experimental Example 12 and Experimental Example 13, a diffusion alloy in which Y was partially added was used. The composition of the shell phase was not inferior to that of the shell compositions of Experimental Example 9 and Experimental Example 11, and the coercive force improvement rate was sufficient.

It was confirmed that the samples prepared in the above experimental examples had a concentration gradient in the core phase and the shell phase according to the element species. In particular, when Si is contained as the M element, the gradient of Si concentration in the core phase is represented by a position-dependent linear function. Sm is often constant within the shell phase. This may be due to the fact that the concentration gradient in the core phase is solid-phase diffusion and the concentration formation in the shell phase is dissolution / reprecipitation, which is mainly caused by the application of the present disclosure technology. It is one of the characteristic organizational forms.

For example, FIGS. 1A and 2A show the results of line analysis of the tissue and composition of Experimental Example 9, and FIGS. 1B and 2B show the results of line analysis of the structure and composition of Experimental Example 11. In addition, since there is a part where both the core phase and the shell phase interfere with each other during the measurement, that part is deleted (the part indicated by "//"). It has been shown that Si is also diffused within the core phase. It was found that Al and Ga also diffuse in the core phase, but not as much as Si. The gradient of Si concentration within the core phase shows a position-dependent linear functional change. Although it is not clear, it may be a linear function in the range where Al and Ga are also diffused. We observed that Co tends to be slightly richer in the shell phase than in the core phase. This means that the amount of Co substitution in the core phase can be kept slightly less than 0.2, which maximizes the magnetic property value, and Co can be concentrated in the shell phase to y to 0.2. To do. By doing so, it is possible to reduce the amount of expensive Co used. Due to the concentration gradient of the elements between the core phase and the shell phase, the lattice constant also changes continuously and the lattice mismatch is eliminated.

On the other hand, in <Step A>, the alloy composition was designed, for example, w = 9, x = 0.60 so that the composition of the entire alloy was Y-rich, and a densified bulk body was obtained by the same method. It was found that inside the bulk body, a shell having a higher Sm concentration than inside the magnetic powder was formed on the outer peripheral portion of the magnetic powder. The composition of the core is Y _0.40 Sm _0.60 (Fe _0.82 Co _0.18 ) _11.07 Al _0.04 Si _0.02 Ti _0.62 Cu _0.25 , whereas the shell The composition was Y _0.29 Sm _0.71 (Fe _0.83 Co _0.17 ) _11.08 Al _0.04 Si _0.02 Ti _0.62 Cu _0.24 .

In order to increase the coercive force of a magnet having a ThMn _12- type crystal structure as the main phase, it is necessary to grasp the crystal structure of the main phase crystal grains in detail. Therefore, a nuclear density analysis was performed using the maximum entropy method (MEM) by powder neutron Rietveld analysis. Since X-rays are diffracted by the spatially spread electron density, the resolution may be inferior in order to determine the position of trace elements with high accuracy. On the other hand, since neutron rays are scattered from the nucleus near the mass point, they are suitable for determining the positions of trace elements with high accuracy. The measurement was carried out at the BL20 structure in the material and life science experimental facility (Material and Life Science Experiment Facility, MLF) of the Japan Proton Accelerator Research Complex (J-PARC). The device (iMATERIA) was used. Measured at room temperature in single frame mode. A hollow vanadium cylinder (φ5.8) is placed in the vanadium container (φ6) and a powder sample (<75 μm) is placed in the gap between the samples so that the absorption can be corrected in the sample containing Sm with large neutron absorption. Was filled (sample shape with cylindrical symmetry). As the analysis data, the data of the back bank after subtracting the background of the vanadium container was used, and the analysis was performed in the range of 0.49Å <d <2.57Å. The analysis codes used were Z-Rietveld 1.0.2 and Z-MEM 0.8.3.

The sample prepared in step A was pulverized to 75 μm or less and evaluated. This sample is a ThMn type ₁₂ compound in which the amount of substitution of the structural stabilizing element Ti is reduced by partial substitution of Y. It was also confirmed that the following occurred regardless of the presence or absence of Cu addition.

FIG. 3 shows the kernel density distribution of the 3D, (100) plane, and (110) plane derived by MEM. The non-isotropic nuclear density distribution corresponds to the anisotropic thermal vibration. It was clarified that there are two types of sites with high nuclear density distribution in the [001] direction of the 2a site, which is a normal rare earth site. In addition, a very small amount of 4g ₁ site, which is an irregular part of the remnants of pseudo-tetragonal crystals, was also observed. Crystallographically, it corresponds to 2b sites (sites containing invading elements) and 4e sites. Since the 2a site is located in the middle between the 4e sites, the 4e site can be seen as a crystal structure forming a dumbbell and irregularly replacing the 2a site. It is presumed that the extruded rare earth element is placed at the 2b site at the place where the rare earth element at the 2a site is replaced with the dumbbell at the 4e site. When this was incorporated into the crystal structure model and Rietveld analysis was performed again, the analysis accuracy of the fitting was improved. As a result, it was found that this irregular replacement rate was about 2%, which was a certain ratio at one site per 25 unit cells. Considering the error and the change according to the amount of Ti, it may be contained up to about 5%. It is possible that this irregular substitution of about 2% to 5% reduces the coercive force. Concentrating the structural stabilizing element at the outer periphery of the crystal to remove such irregular substitutions inside the structure may be equivalent to removing the origin of the magnetization reversal. In fact, in the shells enriched with Ti and Si in Experimental Examples 2 to 11, there is a possibility that the coercive force is improved by reducing or eliminating the ratio of this irregular substitution.

FIG. 4 shows the correspondence of each site when the structure changes from TbCu ₇ type to ThMn ₁₂ type, which is seen in Y—Fe compounds. In the continuous lattice deformation of the Y—Fe-based compound described in Patent Document 3, the deformation in the shortest path from TbCu ₇ type to ThMn ₁₂ type is taught. On the other hand, the structure near the Ti _0.5 composition, which is a transient region from the non-equilibrium phase to the equilibrium phase, which was newly discovered through diligent research, has an irregular substitution part in the direction of _cttra and is bypassed. It was clarified that the structural changes were observed. That is, it was shown that the irregular portion in the _ortho direction inherent in the TbCu ₇ type disappears, and the irregular portion in the c _terra direction is generated and disappears, and is transformed into the ThMn ₁₂ type. Since the (100) in-plane direction of the 2b site is too narrow for rare earth elements to enter, it is not excluded that there is a possibility that other than rare earth elements are contained.

An embodiment of the present disclosure provides a rare earth magnet having main phase crystal grains of a ThMn _12- type compound, which has an improved coercive force while suppressing a decrease in magnetization. Since such rare earth magnets can be suitably used for motors, actuators, and the like, they have various industrial uses.

Claims (5)

A rare earth magnet having a main phase crystal grain of an Fe-based ThMn _12- type compound containing Sm and Y and a grain boundary phase containing Cu.
The main phase crystal grains have a ThMn ₁₂ type core phase and a ThMn ₁₂ type shell phase.
A rare earth magnet in which the Sm concentration in the shell phase is higher than the Sm concentration in the core phase. The main phase crystal grains contain at least one M element selected from the group consisting of Ti, Si, Al, and Ga, and the concentration of the M element in the shell phase is the concentration of the M element in the core phase. The rare earth magnet according to claim 1, which is higher than the above. The rare earth magnet according to claim 2, which contains Si as an M element, and the gradient of the Si concentration at least in the core phase is represented by a position-dependent linear function. In the composition formula _{R1 1-x R2 x (Fe} 1-y Co y) w Ti z Cu α, R1 has at least Y, may further have a Gd, R2 has at least Sm, further It is at least one selected from the group consisting of La, Ce, Nd and Pr, and x, y, z and w are 0.5 ≦ x ≦ 1.0 and 0 ≦ y ≦ 0.4, 8 respectively. A step of preparing a bulk body of an alloy for rare earth magnets, wherein ≦ w ≦ 12, 0.42 ≦ z <0.70, 0.40 ≦ α ≦ 0.70.
A step of preparing a diffusion alloy, which is a composition formula R1 _1-χ R2 _χ Cu _ω- based diffusion alloy and has χ> x, 2 ≦ ω ≦ 4.
By heating the bulk body in contact with the diffusion alloy and diffusing the component elements of the diffusion alloy into the bulk body, the Sm concentration in the shell phase is higher than the Sm concentration in the core phase. ThMn ₁₂ The process of forming a rare earth magnet having main phase crystal grains of a type compound,
A method for manufacturing a rare earth magnet, including. The diffusion alloy is represented by the composition formula R1 _1-χ R2 _χ (Cu _1-υ−τ X _υ Si _τ Zn _σ ) _ω , where X is at least one of Al or Ga, and χ> x, 0 < The method for producing a rare earth magnet according to claim 4, wherein υ ≦ 0.3, 0 <τ <0.5, 0 <σ <0.2, 2 ≦ ω ≦ 4.