JP7187791B2 - Alloys for rare earth magnets - Google Patents

Description

The present invention relates to an alloy for rare earth magnets.

In recent years, there has been a demand for the development of magnets with a reduced content of rare earth elements. As used herein, rare earth elements refer to at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanides. Here, the lanthanoid is a general term for 15 elements from lanthanum (La) to lutetium (Lu). 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 a ferromagnetic alloy containing a relatively small composition ratio of rare earth elements. Are known. Although RT ₁₂ has high magnetization, it suffers from thermal instability of the crystal structure.

Patent Document 1 discloses a rare earth permanent magnet in which a portion of Fe, which is a T element, is partially replaced with Ti, which is a structure stabilizing element, to improve thermal stability in exchange for high magnetization. there is

In Patent Document 2, by partially substituting the R element of the RFe ₁₂ -based compound with the substituting element M1 such as Zr and Hf, the amount of the substituting element M2 such as Ti substituting the transition metal element is reduced and saturated. Rare earth permanent magnets with stabilized _ThMn12 structures are disclosed while retaining their magnetization.

Further, Patent Document 3 discloses an R'--Fe--Co system ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R'--Fe--Co system ferromagnetic It is described that the alloy exhibits high magnetic properties due to having a ThMn ₁₂ -type crystal structure produced by an ultraquenching method.

Further, Patent Document 4 describes that by adding Cu, a non-magnetic and low melting point 1-4 composition (SmCu ₄ phase) phase is generated, and sintering and high coercive force are possible. .

Further, in Patent Document 5, at least one of Sm ₅ Fe ₁₇ -based phase, SmCo ₅ -based phase, Sm ₂ O ₃ -based phase, and Sm ₇ Cu ₃ -based phase is added as a secondary phase to a ThMn ₁₂ -type main phase. It is described that the coercive force can be increased by including the element.

Further, Patent Document 6 describes that the addition of Cu generates a liquid phase and enables the formation of a dense bulk body.

JP-A-64-76703 JP-A-4-322406 JP 2015-156436 A Japanese Patent Application Laid-Open No. 2001-189206 Japanese Unexamined Patent Application Publication No. 2017-112300 WO2016/162990

The single-crystal-like main phase particles used in anisotropic magnets, which are often used as high-performance magnets, are obtained by efficiently pulverizing the raw material alloy (material to be pulverized) into single-crystal units during fine pulverization. . Furthermore, considering the general treatment temperature during the sintering process, the main phase compound is also required to exist stably at least at 900° C. or higher, preferably at 1000° C. or higher.

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

On the other hand, in the rare earth permanent magnet described in Patent Document 2, the _ThMn12 structure is stabilized by substituting a transition metal element with Ti or the like, but the effect is not necessarily sufficient.

The R'--Fe--Co ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structure stabilizing element M (such as Ti), so it exhibits high magnetization, large magnetic anisotropy, and a high Curie temperature. However, due to the non-equilibrium phase, the main phase compound may decompose in a high temperature densification process such as sintering.

The rare earth magnet described in Patent Document 4 may not have high magnetic physical properties due to the large amount of Ti added. Furthermore, there is a concern that the saturation magnetization and magnetic anisotropy may decrease due to solid solution of Cu in the main phase.

In the rare earth magnet described in Patent Document 5, when the rare earth-rich secondary phase Sm ₇ Cu ₃ is used, there is concern that the equilibrium state will shift to a composition richer in rare earth than the main phase during heat treatment, and the main phase ratio will decrease. be done.

In the rare earth magnet described in Patent Document 6, since the Fe element is not replaced with the structure stabilizing element M, high magnetization, large magnetic anisotropy, and a high Curie temperature can be obtained, and the density as a bulk body is high. Since it is a non-equilibrium phase, the main phase compound may decompose in a high temperature process such as sintering at 1000° C. or higher. Furthermore, it is difficult to obtain an anisotropic sintered magnetic powder.

Accordingly, an object of the present disclosure is to provide a rare earth magnet alloy suitable for obtaining anisotropic sintered magnetic powder.

The rare earth magnet alloy of the present disclosure is a rare earth magnet alloy having a main phase and one or more subphases, and the composition of the entire alloy is represented by the following composition formula (1),
R( _Fe , _Co ) _{wzTizCuαSnβ (} 1)
R is at least one rare earth element, and w, z, and α are 8≦w≦13, 0.42≦z<0.70, 0.35≦α≦0.82, and 0<β, respectively ≦0.10 is satisfied.

In certain embodiments, said R consists of R1 and R2,
The overall composition is represented by the following compositional formula (2),
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α Sn _β (2)
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, Sm is 50 mol % or more of the entire R2, and x and y respectively satisfy 0.5≦x≦1.0 and 0≦y≦0.4.

In one embodiment, the main phase has a ThMn ₁₂ -type crystal structure,
The composition of the main phase is represented by the following compositional formula (3),
R1 _1-x' R2 _x' (Fe _1-y' Co _y' ) _12-z'-α' Ti _z' Cu _α' (3)
x', y', z', and α' are respectively 0.5≤x'≤1.0, 0≤y'≤0.4, 0.48≤z'<0.91, and 0 .15≤α'≤0.30.
In some embodiments, z' satisfies 0.48≦z'<0.74.

In one embodiment, the main phase is a phase having a ThMn ₁₂ -type crystal structure, and the secondary phase is mainly composed of a Cu-composition crystal phase and a Sn-based crystal phase, each of which accounts for 50 mol % or more of the total secondary phase.

In one embodiment, the subphase comprises at least a Cu-based KHg ₂ -type crystal structure phase or a Cu-based phase having a composition of rare earth elements and 3d transition elements in a molar ratio of 1:4, and at least Sn It contains a phase of the NdPtSb type crystal structure of the base.

In one embodiment, the KHg type ₂ crystal structure phase is 50% or more by volume of said minor phase.

In one embodiment, the subphase comprises R atoms, and the R atoms present in the subphase have a molar ratio of [R2]/([R1]+[R2]) higher than the composition of the entire alloy.

In one embodiment, the composition formula ratio of the Cu-based subphase is 9 mol % or more and 27 mol % or less, and the Sn-based phase is less than the Cu-based phase

According to the embodiments of the present invention, anisotropic magnetic powder with improved magnetic properties and thermal stability can be obtained with high efficiency.

FIG. 2 is a diagram showing observation results of a polarizing microscope cross-sectional structure and composition analysis results after heat treatment at 1050° C. for 20 minutes for the rare earth magnet alloy of Example 1 in the embodiment of the present disclosure. FIG. 2 is a backscattered electron (BSE) image obtained with a scanning electron microscope (SEM) after heat treatment at 1050° C. for 20 minutes for the rare earth magnet alloy of Example 1 in the embodiment of the present disclosure;

[Composition of alloy for rare earth magnet]
The rare earth magnet alloy of the present disclosure, in an exemplary, non-limiting embodiment, has a primary phase and a secondary phase, and the overall composition is represented by compositional formula (1) below.
R( _Fe , _Co ) _{wzTizCuαSnβ (} 1)
Here, R is at least one rare earth element. Also, w, z, and α satisfy 8≦w≦13, 0.42≦z<0.70, 0.35≦α≦0.82, and 0<β≦0.10, respectively.

In one embodiment, the overall composition is represented by the following compositional formula (2).
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α Sn _β (2)
Here, R is composed of R1 and R2. R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, Sm is 50 mol% or more of the entire R2. R1 is preferably Y only (excluding unavoidable impurities), and R2 is preferably Sm only (excluding unavoidable impurities). x, y, z, and w are respectively 0.5≦x≦1.0, 0≦y≦0.4, 8≦w≦12, 0.42≦z<0.70, 0.35≦α ≦0.82 and 0<β≦0.10 are satisfied.

By including Sm in the rare earth element R, the magnetic anisotropy of the main phase, which is important for increasing the coercive force, can be improved.

As a result of intensive research by the present inventors, as shown in the above composition formulas (1) and (2), by adding Cu to the raw material alloy, the molten metal of the raw material alloy is rapidly cooled and solidified into the alloy. It was found that a rare earth-rich phase (subphase) coexisting with the main phase (hard magnetic phase with high magnetization and magnetic anisotropy) was formed. It was found that the formation of the secondary phase richer in rare earth elements than the main phase facilitates crystal growth of the main phase by heat treatment performed on the rapidly solidified alloy. This heat treatment also makes it possible to easily reduce heterogeneous phases during melting and solidification of the material alloy. Furthermore, the absorption and release of hydrogen by the rare earth-rich subphase causes cracks to occur between the main phase and the subphase or in the subphase, enabling efficient pulverization into single crystal units. These are extremely beneficial in obtaining anisotropic sintered magnetic powder, and can enable mass production of highly oriented anisotropic sintered magnetic powder.

The amounts of R1, R2 and Ti affect the magnetic properties and high temperature stability of the main phase. From the viewpoint of magnetic anisotropy, R2 is desirably half or more of R1 (half or more of the total of R), and the preferred range of x is 0.5≤x≤1.0. From the viewpoint of saturation magnetization, Ti should preferably be as small as possible, but from the viewpoint of high-temperature stability, it should be large. A range of 0.42≦z<0.70 is suitable. In particular, when z is 0.70 or more, the saturation magnetization, magnetic anisotropic magnetic field and Curie temperature all decrease. Note that 50 mol % or less of Ti may be replaced with tungsten (W), vanadium (V), or the like.

From the viewpoint of increasing the magnetic moment and improving the magnetization and magnetic anisotropy at practical temperatures accompanying the increase in the Curie temperature, it is preferable to replace part of Fe with Co. However, if the amount of substitution with Co is too large, the magnetization and magnetic anisotropy will rather be lowered. Specifically, the Co substitution amount y is desirably 0≤y≤0.4, and more desirably 0.1≤y≤0.3.

The amount of Cu is set so that the amount of the generated subphase is an appropriate value. If the amount of the subphase is small, not only is it impossible to eliminate the heterogeneous phase during melting and solidification of the raw material alloy, but it is also difficult to grow crystals to a size sufficient to obtain anisotropic sintered magnetic powder. Also, when the amount of the sub-phase is large, the ratio of the main phase is lowered, so the magnetization of the magnet body is lowered. According to the inventor's experiments, the appropriate amount of Cu is in the range of 0.35≦α≦0.82.

The amount of Sn is important because it determines the amount of Cu dissolved in the main phase to be produced. The addition of Sn produces at least a phase having a NdPtSb type crystal structure (1-1-1 phase), and changes the distribution of Cu. If the amount of Sn added is small, the solid solution amount of Cu in the main phase increases, which is not preferable in terms of saturation magnetization and magnetic anisotropy. It is not easy to inhibit crystal growth and reduce heterogeneous phases. Therefore, the range of 0<β≦0.10 is appropriate.

50 mol % or more of the entire subphase is composed of a crystal phase having a Cu composition (that is, a Cu group richer in R than the main phase) and a Sn group. In some embodiments, the subphase Cu-based phase is predominantly a KHg ₂ -type crystal structure phase (hereinafter 1-2 phase). The subphase may also include a phase having a composition in which the ratio of "R" to "Cu, Fe, and/or Co" is 1:4 (hereinafter referred to as 1-4 composition). On the other hand, the Sn-based phase mainly includes a phase of NdPtSb type crystal structure. Regarding the R element constituting the subphase, both phases have a higher molar ratio of [R2]/([R1]+[R2]) than the composition of the entire alloy. In addition, Fe and Co may slightly dissolve in the subphase. A very small amount of Ti dissolves in the Sn-based phase.

An appropriate amount of w varies depending on the amount of Cu added to the raw material alloy, but 8≤w≤13. If w is too large, a soft magnetic α-(Fe, Co, Ti) phase is produced. If w is too small, 2-17 phases and 3-29 phases are generated. None of these phases are preferable for obtaining magnets with high magnetic properties.

The main phase of the RT _w Ti _z Cu _α Sn _β rare earth magnet alloy thus obtained has a ThMn ₁₂ -type crystal structure in an embodiment. The TnMn ₁₂ -type compound phase in the alloy in the present disclosure can typically exist stably even at 1000° C. or higher. As such, the alloy embodiments of the present disclosure can be suitably used to employ high performance magnet fabrication processes such as sintering.

In general, the "ThMn type ₁₂ crystal structure" is a tetragonal crystal, but in the present invention, the crystal lattice of the tetragonal crystal is slightly distorted to have orthorhombic symmetry, and the atoms in the crystal Even if the periodicity of is slightly disturbed, it is regarded as "ThMn type ₁₂ crystal structure".

Since the main phase to be generated contains Cu, the magnetic property values are different when compared with the main phase not containing Cu, even with the same amount of Ti substitution. However, by adding Sn during melting, the solid solution amount of Cu in the main phase can be suppressed. First, the saturation magnetization decreases at least by the solid solution amount of Cu and Ti. The magnetic anisotropic magnetic field exhibits complex behavior due to the co-substitution of Cu and Ti. Specifically, contrary to the general characteristics of ThMn ₁₂ -type crystal structure compounds in which Cu is not substituted, the magnetic anisotropy field tends to decrease with the addition of Ti. Therefore, it is preferable that both Cu and Ti be as small as possible from the viewpoint of magnetic property values. However, from the viewpoint of high-temperature stability, a large amount of Ti is desirable. Specifically, when the main phase is represented by the composition formula R1 _1-x' R2 _x' (Fe _1-y' Co _y' ) _12-z'-α' Ti _z' Cu _α' , 0.5 ≤ x′≦1.0, 0≦y′≦0.4, 0.48≦z′<0.91, and 0.15≦α′≦0.30 are suitable, and the substitution amount of Cu and Ti is More preferably, 0.48≦z′<0.74 and 0.15≦α′≦0.30.

[Method for producing alloy for rare earth magnet]
<Step A> Melting and solidifying step As a method for producing the R-Fe-Co-Ti-Cu-based rare earth magnet alloy, known methods such as mold casting, centrifugal casting, strip casting, and liquid ultra-quenching can be used. method can be adopted. These methods involve making a molten alloy and then cooling and solidifying the molten alloy. During solidification of the molten alloy, it is desirable to minimize the formation of phases (heterogeneous phases) such as α-(Fe, Co, Ti) phases, which are particularly unfavorable as raw material alloys for magnets. By adopting a method of supplying a molten metal onto a rotating roll and solidifying it to produce an alloy in the form of ribbons or flakes, such as a strip casting method or a liquid ultra-quenching method, which has a relatively high cooling rate, such heterophases generation can be suppressed. When the cooling rate during solidification is low, the grain size of the precipitated heterogeneous phase becomes large. When the grain size of the heterogeneous phase contained in the alloy becomes large, it is difficult for the heterophase to disappear in the subsequent heat treatment step B.

When a molten alloy is quenched and solidified at a high cooling rate as in the liquid ultraquenching method, "nanocrystals" with nanometer-order sizes are produced in the solidified alloy. Anisotropic magnetic powder cannot be obtained by pulverizing the solidified alloy in the state of "nanocrystals". However, even nanocrystals can be easily grown into crystal grains of 10 μm or more suitable for obtaining anisotropic magnetic powder by going through the subsequent heat treatment step B.

<Step B> Heat Treatment Step By applying heat treatment to the alloy of the present invention, the following can be achieved.
(1) It reduces the amount of foreign phases produced during the solidification process.
(2) Coarse crystal grains. This is effective for easily obtaining a powder composed of single-crystal-like particles, which is useful as a raw material for an anisotropic sintered magnet, by a pulverization method.

The melting point of the 1-2 phase is around 860.degree. C., and the melting point of the 1-4 composition phase is around 880.degree. The melting point of the Sn group 1-1-1 phase is 1000° C. or higher. Since the Sn-based phase hardly dissolves the main phase, it does not contribute to the crystal growth of the main phase. Therefore, the heat treatment temperature is preferably 900° C. or higher and 1250° C. or lower, more preferably 1000° C. or higher and 1100° C. or lower.

The heat treatment time depends on the heat treatment temperature, but is preferably from 5 minutes to 50 hours. If the time is too short, there will be insufficient reaction to eliminate the heterogeneous phase or insufficient grain growth. If the time is too long, evaporation and oxidation of the rare earth elements will occur, and the operational efficiency will be poor. At this heat treatment temperature, at least part of the subphase becomes a liquid phase, and part of the main phase dissolves and reprecipitates. For this reason, the crystal grains grow dramatically in the main phase as compared with the case where the liquid phase is not generated. In addition, heterogeneous phases generated during rapid solidification of the molten alloy can be easily eliminated if the grain size is not large.

The amount of the subphase in the alloy before the step B is desirably 9 mol % or more and 27 mol % or less. If the amount of the subphase is small, not only will the heterogeneous phase in the alloy not disappear even by the heat treatment in step B, but it will also be difficult to grow the crystals to a size sufficient to obtain an anisotropic sintered magnetic powder. Also, when the amount of the sub-phase is large, the ratio of the main phase is lowered, so the magnetization of the magnet body is lowered. The alloy thus obtained can be used as a rare earth magnet alloy that can be efficiently pulverized into single crystal units by hydrogen embrittlement.

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

[Example 1]
<Process A>
Raw material metals of Y, Sm, Fe, Co, Ti, Cu, and Sn with a purity of 99.9% or more were weighed in increments, taking into account the evaporation of rare earth elements during melting. After sufficiently melting these raw metals in the tapping pipe of a liquid ultraquenching device (melt spinning device) to form a molten alloy, the molten metal is tapped on a Cu roll rotating at a roll peripheral speed of 15 m/s. did. The molten metal contacted the surface of a roll rotating at high speed, was rapidly removed from the heat, extended in the form of a ribbon, and solidified. In Example 1 and Comparative Example 2, the molten metal was tapped on a Cu roll rotating at 15 m/s, and in Comparative Example 1, the molten metal was tapped on a Cu roll rotating at 40 m/s. In this way, ultraquenched ribbons having the compositions shown in Table 1 were produced.

<Process B>
The ultraquenched ribbon produced in step A was wrapped in Nb foil and heat-treated at 1050° C. for 20 minutes in an Ar stream.

Tables 1 and 2 show the composition of the alloy, the composition of the main phase identified by SEM-EDX analysis, and the saturation magnetization Ms (T), magnetic anisotropy field Ha (T) and Curie temperature Tc (° C.) at room temperature. indicates However, a vibrating sample magnetometer capable of applying a magnetic field up to 10 T was used to evaluate saturation magnetization and magnetic anisotropic magnetic field. Since the magnetic particles are isotropic, the saturation magnetization was identified using the square-law saturation asymptotic law, and the magnetic anisotropy field was identified using the singularity detection method. The Curie temperature was defined as the inflection point of temperature change of magnetization using a thermomagnetic balance.

As shown in Table 2, the main phase of Comparative Example 1 has high saturation magnetization and magnetic anisotropic magnetic field. However, since Cu is not added as described above, a phase rich in rare earth elements coexisting with the main phase is not generated, and crystal growth of the main phase by heat treatment becomes difficult. Furthermore, cracks cannot be generated between the main phase and the subphase, and efficient pulverization into single crystal units cannot be performed. Therefore, an anisotropic sintered magnet powder capable of high orientation cannot be produced. Also, by adding Cu, an anisotropic sintered magnet powder capable of high orientation can be produced, but as shown in Comparative Example 2 in Table 2, the saturation magnetization and magnetic anisotropic magnetic field of the main phase are lowered. end up On the other hand, in Example 1 of the present invention, compared with Comparative Example 1 in which Sn was not added, the amount of Cu solid solution in the main phase was reduced by adding Sn, and the main phase had high saturation magnetization and high magnetism. It is possible to have an anisotropic magnetic field.

FIG. 1 shows the results of observing the cross-sectional structure of the rare earth magnet alloy of Example 1 with a polarizing microscope after heat treatment at 1050° C. for 20 minutes. FIG. 2 shows a backscattered electron (BSE) image obtained with a scanning electron microscope (SEM) after heat treatment at 1050° C. for 20 minutes for the rare earth magnet alloy of Example 1. FIG. Table 3 shows the results of composition analysis by SEM-EDX.

From this result, it can be seen that the heat treatment causes microcrystals on the order of hundreds of nanometers to grow into coarse crystal grains of 10 μm or more. It can also be seen that the subphases contain a Cu-based 1-2 phase, a 1-4 composition phase, and a Sn-based 1-1-1 phase. At least the above three subphases have compositions that are Sm richer than the Y/Sm ratio of the entire alloy. Powder X-ray diffraction confirmed that this 1-1-1 phase was of the NdPtSb type.

[Examples 2 to 5]
<Process A>
Raw material metals of Y, Sm, Fe, Co, Ti, Cu, and Sn with a purity of 99.9% or more are evaporated so as to have the compositions shown in Comparative Examples 3 to 5 and Examples 2 to 5 when the rare earth elements are melted. was added and weighed in increments. After being sufficiently melted in the tapping pipe, the molten metal was tapped on a Cu roll rotating at 15 m/s to prepare a super-quenched ribbon.

<Process B>
The ultraquenched ribbon produced in step A was wrapped in Nb foil and heat-treated at 1050° C. for 20 minutes in an Ar stream.

Table 4 shows the main phase content, subphase content and alloy composition after heat treatment, the intrinsic content of Cu in the main phase after heat treatment, and the presence or absence of different phases after heat treatment. Examples 2 to 5 of the present invention all have smaller amounts of Cu in the main phase than Comparative Examples 3 to 5. It can be seen that Comparative Examples 3 to 5 do not have high volume magnetization and magnetic anisotropic magnetic field because the amount of Cu solid solution in the main phase is large. Comparing Example 2 and Example 3, the amount of Cu solid solution in the main phase is small in both cases, and the volume magnetization and magnetic anisotropy field are large, but in Example 2, the Sn-based phase is more than half the Cu-based phase. Therefore, the heteromorphic phase cannot be eliminated. This can also be seen by comparing Example 4 and Example 5, which have a large amount of subphase. Therefore, the Sn-based phase is preferably less than half of the Cu-based phase.

Since the rare earth magnet alloy of the present disclosure contains a main phase with improved magnetic properties and thermal stability and a rare earth-rich sub phase, it can be suitably used to produce anisotropic sintered magnetic powder. Anisotropic sintered magnetic powder can be suitably used for producing sintered magnets. Sintered magnets are used in various motors and actuators, and have various industrial uses.

Claims (9)

A rare earth magnet alloy having a main phase and one or more subphases, wherein the composition of the entire alloy is represented by the following compositional formula (1),
R( _Fe , _Co ) _{wzTizCuαSnβ (} 1)
R is at least one rare earth element,
w, z, and α are each 8≤w≤13,
0.42≦z<0.70,
0.35≦α≦0.82, and 0<β≦0.10
An alloy for rare earth magnets that satisfies said R is composed of R1 and R2,
The overall composition is represented by the following compositional formula (2),
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α Sn _β (2)
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1,
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is 50 mol% or more of the total of R2,
x and y are each
0.5≦x≦1.0,
0≤y≤0.4,
The rare earth magnet alloy according to claim 1, which satisfies The main phase has a ThMn ₁₂ -type crystal structure,
The composition of the main phase is represented by the following compositional formula (3),
R1 _1-x' R2 x' (Fe _1-y' Co _y' ) _{12-z'-α' Ti z} ' Cu _α' (3)
x', y', z', and α' are each
0.5≦x′≦1.0,
0≦y′≦0.4,
0.48≦z′<0.91, and
3. The rare earth magnet alloy according to claim 2, which satisfies 0.15≤α'≤0.30. z′, is 0.48≦z′<0.74
The rare earth magnet alloy according to claim 3, which satisfies The main phase is a phase having a ThMn ₁₂ type crystal structure,
5. The alloy for a rare earth magnet according to claim 3, wherein 50 mol % or more of said subphase is composed mainly of a Cu-composition crystal phase and a Sn-based crystal phase. The subphase includes at least a Cu-based KHg ₂ type crystal structure phase or a Cu-based phase having a composition of a rare earth element and a 3d transition element in a molar ratio of 1:4, and at least a Sn-based NdPtSb type 6. The rare earth magnet alloy according to claim 5, comprising a crystalline phase. 7. The rare earth magnet alloy according to claim 6, wherein the KHg type ₂ crystal structure phase accounts for 50% or more of said subphase by volume. 8. Any one of claims 5 to 7, wherein the subphase contains R atoms, and the R atoms present in the subphase have a molar ratio of [R2]/([R1]+[R2]) higher than the composition of the entire alloy. The alloy for rare earth magnets according to 1. 9. The rare earth magnet alloy according to claim 5, wherein the composition formula ratio of the Cu-based subphase is 9 mol % or more and 27 mol % or less, and the Sn-based phase is less than the Cu-based phase.

Images