JP2019044259A - Alloy for rare earth magnet - Google Patents

Abstract

To provide an anisotropic magnetic powder of which magnetic property and heat stability are enhanced at high performance.SOLUTION: The disclosed alloy for rare earth magnet has a composition represented by R(Fe,Co)TiCu, wherein R is at least one kind of rare earth elements, and w, z and α satisfy 8≤w≤12, 0.42≤z<0.70 and 0.40≤α≤0.70.SELECTED DRAWING: Figure 1B

Description

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

In recent years, there has been a demand for the development of a magnet with a reduced content of rare earth elements. In this specification, the rare earth element refers to at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanoid. Here, the lanthanoid is a generic name of 15 elements from lanthanum (La) to lutetium (Lu). As a ferromagnetic alloy in which the composition ratio of the rare earth element 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. Are known. RT ₁₂ has high magnetization, but has a problem that its crystal structure is thermally unstable.

  Patent Document 1 discloses a rare earth permanent magnet in which part of Fe as a T element is partially replaced by Ti as a structure stabilizing element, and the thermal stability is improved in exchange for high magnetization. Yes.

In Patent Document 2, the R element of the RFe _12- based compound is partially substituted with a substitution element M1 such as Zr or Hf, so that the amount of substitution element M2 such as Ti that substitutes the transition metal element is reduced and saturated. A rare earth permanent magnet is disclosed in which the ThMn ₁₂ structure is stabilized while maintaining the magnetization.

Patent Document 3 discloses an R′—Fe—Co based ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R′—Fe—Co based ferromagnetic alloy is disclosed. It is described that the alloy exhibits a high magnetic property because it has a ThMn ₁₂ type crystal structure formed by a rapid quenching method.

Patent Document 4 describes that addition of Cu produces a non-magnetic and low melting point 1-4 composition (SmCu ₄ phase), which enables sintering and high coercivity. .

Patent Document 5 discloses at least one of an Sm ₅ Fe ₁₇ phase, an SmCo ₅ phase, an Sm ₂ O ₃ phase, and an Sm ₇ Cu ₃ phase as a subphase with respect to a ThMn ₁₂ type main phase. It is described that inclusion thereof can increase the coercive force.

  Patent Document 6 describes that by adding Cu, a liquid phase is generated and a dense bulk body can be formed.

JP-A 64-76703 JP-A-4-322406 Japanese Patent Laying-Open No. 2015-156436 JP 2001-189206 A JP 2017-112300 A International Publication No. 2016/162990

  Single-crystal-like main phase particles used in anisotropic magnets that are widely used as high-performance magnets are obtained by high-efficiency pulverization of raw material alloys (substances to be crushed) to single-crystal units. . Furthermore, in consideration of a general processing temperature in the sintering step, the main phase compound is also required to be stably present at least 900 ° C. or more, preferably 1000 ° C. or more.

  Although the rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to elemental substitution of Fe by Ti, since the amount of Fe substitution by Ti is large, magnetization is reduced by that amount, 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, although the ThMn ₁₂ structure is stabilized by substituting the transition metal element with Ti or the like, the effect is not necessarily sufficient.

  Since the R′-Fe—Co based ferromagnetic alloy described in Patent Document 3 does not substitute the Fe element with the structural stabilizing element M (Ti or the like), it has high magnetization, large magnetic anisotropy, and high Curie temperature. Although it is obtained, since it is a non-equilibrium phase, the main phase compound may decompose in a densification process at a high temperature such as sintering.

  In the rare earth magnet described in Patent Document 4, the magnetic property value may not be high due to the large amount of Ti added.

In the rare earth magnet described in Patent Document 5, 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 rather than the main phase during heat treatment and the main phase ratio decreases. Is done.

  In the rare earth magnet described in Patent Document 6, since the Fe element is not substituted with the structural stabilizing element M, high magnetization, large magnetic anisotropy and high Curie temperature can be obtained, and the density as a bulk body is high. Because of the non-equilibrium phase, the main phase compound may decompose in a process at a high temperature such as sintering at 1000 ° C. or higher. Furthermore, it is difficult to obtain 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 alloy for rare earth magnets of the present disclosure is, in a non-limiting exemplary embodiment, 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) w-} z Ti z Cu α (1)
Here, R is at least one of rare earth elements, and w, z, and α satisfy 8 ≦ w ≦ 13, 0.42 ≦ z <0.70, and 0.40 ≦ α ≦ 0.70, respectively. Satisfied.

In an embodiment, the entire composition is represented by the following composition formula (2).
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α (2)
Here, R is composed of R1 and R2, R1 is Y or Y and Gd, Y is 50 mol% or more of the entire R1, and R2 is selected from the group consisting of Sm, La, Ce, Nd and Pr. At least one kind, which necessarily contains Sm, and Sm is 50 mol% or more of the entire R2. x and y satisfy 0.5 ≦ x ≦ 1.0 and 0 ≦ y ≦ 0.4, respectively.

In one embodiment, the main phase has a ThMn ₁₂ type crystal structure, and the composition of the main phase is represented by the following composition formula (3).
R1 _{1-x ′} R2 _{x ′} (Fe _{1-y ′} Co _{y ′} ) _{12-z′-α ′ Tiz ′} Cu _{α ′} (3)
Here, x ′, y ′, z ′, and α ′ are 0.5 ≦ x ′ ≦ 1.0, 0 ≦ y ′ ≦ 0.4, 0.48 ≦ z ′ <0.91, respectively. And 0.24 ≦ α ′ ≦ 0.37 is satisfied.

  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 subphase is mainly a crystal phase having a Cu composition of 50 mol% or more of the entire subphase.

In one embodiment, the crystal structure of the subphase is KHg type ₂ .

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

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

  In a certain embodiment, the weight ratio of the said subphase in the whole alloy for the said rare earth magnets is 3 wt% or more and 10 wt% or less.

  In certain embodiments, the subphase absorbs and releases hydrogen.

  In certain embodiments, hydrogen absorption and release occurs at a temperature of at least 700 ° C. or lower.

  In one embodiment, cracks are generated between or in the main phase and the sub phase.

  In one embodiment, the maximum value of the X-ray intensity in the powder X-ray diffraction pattern having a lattice plane spacing value of 3.03 to 3.10 mm is a lattice plane spacing value of 2.50 to 2.57. It is larger than the maximum value of the X-ray intensity in the following range, and the α- (Fe, Co, Ti) phase is 10% by weight or less.

  In some embodiments, the α- (Fe, Co, Ti) phase is 5 wt% or less.

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

Y _0.4 Sm _0.6 (Fe _0.80 Co _0.20 ) _8.79 Ti _0.42 Cu _0.67 composition rare earth magnet alloy of Example 6 in the embodiment of the present disclosure at 1050 ° C. for 20 minutes It is a figure which shows the observation result of the cross-sectional structure | tissue by the polarization microscope before performing heat processing of (3). Y _0.4 Sm _0.6 (Fe _0.80 Co _0.20 ) _8.79 T _i0.42 Cu _0.67 composition rare earth magnet alloy of Example 6 in the embodiment of the present disclosure at 1050 ° C. for 20 minutes It is a figure which shows the observation result of the cross-sectional structure | tissue by the polarizing microscope after the heat processing of this. It is a figure which shows the backscattered electron (BSE) image obtained with the scanning electron microscope (SEM) of the sample of FIG. 1B. It is a figure which shows the composition analysis result of each phase in the sample of FIG. 1B. It is a figure which shows the temperature dependence of the volume magnetization of the main phase of Example 3 to 6 in embodiment of this indication. It is a graph which shows the temperature dependence of the magnetic anisotropy magnetic field of the sample of Examples 3 to 6 in the embodiment of the present disclosure. It is a figure which shows the powder X-ray-diffraction pattern before and behind heat processing for 20 minutes of 1050 degreeC of the sample of Example 10 in embodiment of this indication. It is a figure which shows the temperature dependence of the hydrogen discharge amount of Example 13 in embodiment of this indication. It is a figure which shows the powder X-ray-diffraction pattern of the sample of Examples 15-18 in embodiment of this indication. It is a figure which shows the BSE image obtained by SEM of the sample of Example 15 in embodiment of this indication. It is a figure which shows the BSE image obtained by SEM of the sample of Example 16 in embodiment of this indication. It is a figure which shows the BSE image obtained by SEM of the sample of Example 17 in embodiment of this indication. It is a figure which shows the BSE image obtained by SEM of the sample of Example 18 in embodiment of this indication.

[Composition of rare earth magnet alloy]
In an exemplary and non-limiting embodiment, the rare earth magnet alloy of the present disclosure has a main phase and a subphase, and the entire composition is represented by the following composition formula (1).
_{R (Fe, Co) w-} z Ti z Cu α (1)
Here, R is at least one rare earth element. Further, w, z, and α satisfy 8 ≦ w ≦ 13, 0.42 ≦ z <0.70, and 0.40 ≦ α ≦ 0.70, respectively.

In an embodiment, the entire composition is represented by the following composition formula (2).
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α (2)
Here, R is composed of R1 and R2. R1 is Y or Y and Gd, Y is 50 mol% or more of the entire R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd, and Pr, and necessarily includes Sm, Sm is 50 mol% or more of the entire R2. R1 is preferably only Y (excluding inevitable impurities), and R2 is preferably only Sm (excluding inevitable impurities). x and y satisfy 0.5 ≦ x ≦ 1.0 and 0 ≦ y ≦ 0.4, respectively.

  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 alloy of the raw material alloy is rapidly cooled and solidified in the alloy. It was found that a rare earth-rich phase (subphase) coexisting with the main phase (hard magnetic phase having high magnetization and magnetic anisotropy) was formed. It has been found that the generation of a sub-phase rich in rare earth than the main phase makes it possible to easily grow the main phase by heat treatment performed on the rapidly solidified alloy. In addition, this heat treatment makes it possible to easily reduce foreign phases during melting and solidification of the raw material alloy. Furthermore, when the rare earth-rich subphase absorbs and releases hydrogen, cracks occur between the main phase and the subphase or in the subphase, and the single crystal unit can be efficiently pulverized. These are extremely useful in obtaining anisotropic sintered magnetic powder, and can enable mass production of highly sintered 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 than R1 (half or more of the entire R), and a preferable range of x is 0.5 ≦ x ≦ 1.0. Further, Ti is preferably as small as possible from the viewpoint of saturation magnetization, but more is desirable from the viewpoint of high-temperature stability. A range of 0.42 ≦ z <0.70 is appropriate. In particular, when z is 0.70 or more, the saturation magnetization, the magnetic anisotropy magnetic field, and the Curie temperature all decrease. Note that 50 mol% or less of Ti may be substituted with tungsten (W), vanadium (V), or the like.

  Moreover, it is preferable to substitute a part of Fe with Co from the viewpoint of improvement of magnetization at practical temperature and improvement of magnetic anisotropy accompanying increase in magnetic moment and improvement of Curie temperature. However, if the amount of substitution by Co is too large, the magnetization and magnetic anisotropy are reduced. Specifically, the Co substitution amount y is preferably 0 ≦ y ≦ 0.4, and more preferably 0.1 ≦ y ≦ 0.3.

  The amount of Cu is set so that the amount of the subphase to be generated becomes an appropriate value. When the amount of the subphase is small, not only the heterogeneous phase at the time of melting and solidifying the raw material alloy cannot be lost, but also it is not easy to grow the crystal to a size sufficient to obtain anisotropic sintered magnetic powder. Moreover, since the ratio of a main phase will fall when there is much quantity of a subphase, the magnetization as a magnet body will fall. According to the inventor's experiment, the appropriate amount of Cu is in the range of 0.40 ≦ α ≦ 0.70.

The generated subphase is a crystal phase having a Cu composition of 50 mol% or more of the entire subphase (that is, a Cu group that is R richer than the main phase). In an embodiment, the subphase mainly includes a phase having a crystal structure of KHg ₂ type (hereinafter referred to as 1-2 phase). The subphase may also include a phase having a composition (hereinafter referred to as a 1-4 composition) in which the ratio of “R” to “Cu, Fe, and / or Co” is 1: 4. Regarding the R element constituting the subphase, in both phases, the molar ratio [R2] / ([R1] + [R2]) is higher than the composition of the entire alloy. In the subphase, Fe and Co may be slightly dissolved. Ti hardly dissolves in both phases.

  The 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 generated. If w is too small, a 2-17 phase or a 3-29 phase is generated. Neither of these phases is preferred for obtaining a magnet with high magnetic properties.

The main phase of the rare earth magnet alloy thus obtained has a ThMn ₁₂ type crystal structure in the embodiment. The TnMn ₁₂ type compound phase in the alloy in the present disclosure can typically exist stably even at 1000 ° C. or higher. For this reason, the embodiments of the alloy of the present disclosure can be suitably used to adopt a high-performance magnet manufacturing process such as a sintering method.

In general, the “ThMn _12- type crystal structure” is a tetragonal crystal. However, in the present invention, the tetragonal crystal lattice is slightly distorted to have orthorhombic symmetry, and the atoms in the crystal. Even when the periodicity is slightly disturbed, it is regarded as a “ThMn _12- type crystal structure”.

Since the main phase to be produced contains Cu, the magnetic property values are different even with the same Ti substitution amount when compared with the main phase not containing Cu. First, the saturation magnetization is lowered by at least the amount in which Cu and Ti are dissolved. The magnetic anisotropy magnetic field exhibits complicated behavior as co-substitution of Cu and Ti. Specifically, contrary to the general characteristics of a compound of ThMn ₁₂ type crystal structure in which Cu is not substituted, the magnetic anisotropic magnetic field tends to decrease with the addition of Ti. Therefore, both Cu and Ti are preferably as small as possible from the viewpoint of magnetic property values for both elements. However, from the viewpoint of high temperature stability, a larger amount of Ti is desirable. Specifically, when the main phase is expressed by the composition formula R1 _{1-x ′} R2 _{x ′} (Fe _{1−y ′} Co _{y ′} ) _{12−z′−α ′ Tiz ′} Cu _{α ′} , 0.5 ≦ x ′ ≦ 1.0, 0 ≦ y ′ ≦ 0.4, 0.48 ≦ z ′ <0.91, 0.24 ≦ α ′ ≦ 0.37 are appropriate, and the substitution amount of Cu and Ti is More preferably, 0.48 ≦ z ′ <0.74 and 0.24 ≦ α ′ ≦ 0.37.

[Production method of rare earth magnet alloy]
<Step A> Step of melting and solidifying Known methods such as a die casting method, a centrifugal casting method, a strip casting method, and a liquid ultra-quenching method as a method for producing an R-Fe-Co-Ti-Cu rare earth magnet alloy The method can be adopted. In these methods, after preparing a molten alloy, the molten metal is cooled and solidified. It is desirable to suppress as much as possible the generation of an unfavorable phase (heterophase) as a raw material alloy for magnets such as an α- (Fe, Co, Ti) phase during solidification of the molten alloy. By adopting a method of producing a ribbon or flake-like alloy by supplying a molten metal onto a rotating roll and solidifying it, such as a strip casting method or a liquid ultra-cooling method, which has a relatively high cooling rate. Generation can be suppressed. When the cooling rate at the time of solidification is low, the grain size of the different phase that precipitates increases. When the grain size of the heterogeneous phase contained in the alloy becomes large, it is difficult for the heterogeneous phase to disappear in the heat treatment step of the next step B.

  When the molten alloy is rapidly cooled and solidified at a high cooling rate as in the liquid superquenching method, “nanocrystals” having a size on the order of nanometers are formed in the alloy after solidification. In the “nanocrystal” state, anisotropic magnetic powder cannot be obtained even when the solidified alloy is pulverized. However, even nanocrystals can be easily grown into crystal grains having a size of 10 μm or more suitable for obtaining anisotropic magnetic powder by performing the heat treatment step of the subsequent step B.

<Step B> Heat treatment step By applying a heat treatment to the alloy of the present invention, the following can be realized.
(1) Reduce the amount of heterogeneous phase generated during the solidification process.
(2) The crystal grains are coarsened. This is effective for easily obtaining a powder composed of single crystal-like particles useful as a raw material for an anisotropic sintered magnet by a pulverization method.

  Although it depends on the composition of the alloy, the melting point of the 1-2 phase is around 860 ° C., and the melting point of the phase of the 1-4 composition is around 880 ° C. 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 is preferably from 5 minutes to 50 hours, depending on the heat treatment temperature. If the time is too short, the reaction may not be sufficient to eliminate the heterogeneous phase, or the grain growth may be insufficient. If the time is too long, evaporation and oxidation of rare earth elements occur, and the operational efficiency is poor. At this heat treatment temperature, at least a part of the subphase becomes a liquid phase and a part of the main phase is dissolved and reprecipitated. For this reason, crystal grains grow dramatically in the main phase compared to the case where no liquid phase is generated. Also, the heterogeneous phase generated during the rapid solidification of the molten alloy can be easily eliminated if the grain size is not large.

  The weight ratio of the subphase in the alloy before performing Step B is desirably 3 wt% or more and 10 wt% or less of the entire alloy. If the amount of the subphase is small, not only the heterogeneous phase in the alloy cannot be lost even by the heat treatment in Step B, but it is not easy to grow the crystal to a size sufficient to obtain anisotropic sintered magnetic powder. Moreover, since the ratio of a main phase will fall when there is much quantity of a subphase, the magnetization as a magnet body will fall. When the actual subphase amount was identified by X-ray Rietveld analysis of a sample having a subphase weight ratio of 3%, the amount of subphase introduced was the same as the amount of subphase in the alloy.

<Step C> Step of introducing cracks by absorbing and releasing hydrogen Since the subphase contained in the alloy has a rare earth-rich composition containing a relatively large amount of rare earth, it can absorb and release hydrogen. In particular, the subphase can absorb hydrogen significantly by heat treatment in hydrogen. According to this alloy, for example, hydrogen absorption occurs at a temperature of 250 ° C. to 400 ° C., and hydrogen is released between 300 ° C. and 660 ° C. For this reason, the alloy can be heated to 350 ° C. in hydrogen to absorb the hydrogen and then switched to a vacuum atmosphere or an inert gas atmosphere to release hydrogen. In that case, the temperature for switching to a vacuum atmosphere or an inert gas atmosphere is preferably less than 600 ° C. When this alloy is exposed to a hydrogen atmosphere at a temperature of 600 ° C. or higher, the main phase may be decomposed by a hydrogenation-disproportionation reaction to generate a different phase such as an α- (Fe, Co, Ti) phase. The proportion of the α- (Fe, Co, Ti) phase in the alloy after the hydrogen treatment is desirably 10% by weight or less. More preferably, the ratio of the α- (Fe, Co, Ti) phase is 5% by weight or less. If the ratio of the α- (Fe, Co, Ti) phase is high, there is a possibility that the main phase is not sufficiently regenerated by the recombination reaction even after the subsequent heat treatment in a vacuum atmosphere or an inert gas atmosphere. is there. By absorbing and desorbing hydrogen, the rare earth-rich phase (subphase) undergoes volume expansion and contraction, and cracks occur between the main phase crystal grains and between the main phase crystal grains and the subphase. As an example, when the KHg type ₂ phase reacts with hydrogen, the lattice spacing value in the powder X-ray diffraction pattern changes to a phase having a peak at 3.03 to 3.10. As a result, the volume of the subphase increases and cracks occur between the main phase crystal grains. Single crystal units due to cracks between main phase grains and between main phase grains and subphases, which increase the probability of magnetic powder breaking in cracks during the grinding process using a jet mill, stamp mill, ball mill, etc. It is possible to obtain highly-orientable anisotropic magnetic powder containing a large amount of fine powder.

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

[Examples 1 to 6]
<Process A>
Y, Sm, Fe, Co, Ti, and Cu raw metals having a purity of 99.9% or more were weighed in increments in consideration of evaporation of rare earth elements during melting. These raw metals are sufficiently melted in a tapping pipe of a liquid ultra-quenching device (melt spinning device) to form a molten alloy, and then the molten metal is tapped onto a Cu roll rotating at a roll peripheral speed of 15 m / s. did. The molten metal contacted the surface of the roll rotating at a high speed and was quickly removed from the heat, and extended into a ribbon shape and solidified. Thus, ultra-quenched ribbons having the compositions shown in Table 1 were produced. As a comparative example, an ultra-quenched ribbon having an alloy composition Y _0.40 Sm _0.60 (Fe _0.80 Co _0.20 ) _8.33 Ti _0.70 Cu _0.63 was also produced.

<Process B>
The ultra-quenched ribbon produced in step A was included in the Nb foil, and heat treatment was performed at 1050 ° C. for 20 minutes in an Ar stream to produce the alloys of Examples 1 to 6 and Comparative Example 1.

Table 1 shows the composition of the alloy, the composition of the main phase identified by SEM-EDX analysis, the saturation magnetization (Ms) at room temperature, the magnetic anisotropy field (μ ₀ H _a ), and the Curie temperature (T _c ). Indicates. Saturation magnetization (M _s ) and magnetic anisotropic magnetic field (μ ₀ H _a ) were evaluated using a vibrating sample magnetometer (VSM) capable of applying a magnetic field up to 10T. Since the magnetic powder is isotropic, the saturation magnetization was identified using the square law saturation asymptotic law, and the magnetic anisotropy field was identified using the singular point detection method. The Curie temperature (T _c ) was defined as the inflection point of a curve obtained with a thermomagnetic balance. Here, the thermomagnetic balance is a device that applies a magnetic field to a sample portion of a thermobalance (TG) with a permanent magnet and detects a temperature change of magnetic absorption power in the sample as a weight value of TG.

  As shown in Table 1, when the amount of Ti added to the raw material alloy decreases, the Ti content of the main phase in the finally obtained alloy (after Step B) decreases, while the Cu content hardly changes. . This is because Ti does not dissolve in the subphase but only in the main phase.

According to this example, both the saturation magnetization at room temperature and the magnetic anisotropy field tend to increase generally as the Ti addition amount decreases. Furthermore, the Curie temperature also increased as the Ti addition amount decreased. Further, in the comparative example in which the addition amount of Ti was 0.70, the saturation magnetization (M _s ), the magnetic anisotropic magnetic field (H _a ), and the Curie temperature (T _c ) all decreased.

As an example of the structure change before and after the heat treatment in Step B, FIGS. 1A and 1B show Y _0.4 Sm _0.6 (Fe _0.8 Co _0.20 ) _8.79 Ti _0. The observation result of the structure | tissue using the polarization microscope before and behind heat processing of the rare earth magnet alloy of ₄₂ Cu _0.67 is shown. From FIG. 1A and FIG. 1B, it can be seen that microcrystals having a size on the order of several hundreds of nanometers before the heat treatment grow into coarse crystal grains of 10 μm or more by the heat treatment in the step B. In addition, it can be seen that the subphase is clearly more in the 1-2 phase than in the 1-4 composition phase (the black portion of the grain boundary is the 1-2 phase, white (orange in color) linear) And wedge-shaped part is 1-4 phase).

  2A and 2B show the composition analysis results by SEM-EDX. From FIG. 2B, it can be seen that both the 1-2 phase and the 1-4 composition phase of the subphase are Sm richer than the Y / Sm ratio of the whole alloy. Further, the phase having the 1-4 composition contains more Fe and Co inside than the 1-2 phase. Furthermore, it was not possible to measure in both phases where Ti was detectable.

  3A and 3B are graphs showing the temperature dependence (from room temperature to 140 ° C.) of the volume magnetization and magnetic anisotropy field of the main phase of the alloy compositions of Examples 3 to 6, respectively. However, the ratio between the main phase and the subphase was identified by X-ray Rietveld analysis, and the subphase was treated as non-magnetic. Volume magnetization decreases with increasing temperature. As shown in FIG. 3A, Examples 3 to 6 (z ′ in the main phase composition is 0.48 to 0.74) are higher than Comparative Example 1 (z ′ in the main phase composition is 0.91). Is obtained. Moreover, as shown in FIG. 3B, Examples 4 to 6 (z ′ in the main phase composition is 0.48 to 0.67) are more in Example 5 (z ′ is 0.74) and Comparative Example 3 (z The lowering of the magnetic anisotropy magnetic field due to the temperature rise (20 ° C. to 100 ° C., which is practically important) as compared with “0.91” is suppressed. Therefore, preferable ranges are 0.48 ≦ z ′ ≦ 0.67 for the main phase composition and 0.42 ≦ z ≦ 0.50 for the alloy composition.

[Examples 7 to 12]
<Process A>
The raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of the evaporation of rare earth elements during melting. Each element was weighed so that the amounts of main phase and subphase shown in Table 2 were obtained. These raw metals were sufficiently dissolved in the tapping pipe of a liquid ultra-quenching device (melt spinning device), and then the molten metal was tapped onto a Cu roll rotating at a roll peripheral speed of 15 m / s. The molten metal contacted the surface of the roll rotating at a high speed and was quickly removed from the heat, and extended into a ribbon shape and solidified. Thus, ultra-quenched ribbons having the compositions shown in Table 2 were produced.

<Process B>
The ultra-quenched ribbon produced in step A was included in the Nb foil, and was subjected to heat treatment at 1050 ° C. for 20 minutes in an Ar flow to obtain alloys of Examples 7-12.

Table 2 also describes the presence or absence of a heterogeneous phase after heat treatment with respect to the weight ratio of the main phase and the subphase. When the liquid phase amount was small, the heterogeneous phase could not be lost. When the liquid phase amount was 3 wt% or more, the heterogeneous phase disappeared, and a rare earth magnet alloy having only the main phase and the subphase could be produced. FIG. 4 shows the powder X-ray diffraction results of the ultra-quenched ribbon before and after heat treatment with a subphase ratio of 3 wt%. The diffraction peaks of the TbCu ₇ phase and α- (Fe, Co, Ti) phase that existed mainly as different phases before the heat treatment disappeared after the heat treatment at 1050 ° C. for 20 minutes, and the diffraction peaks of the main phase and the subphase Only was observed.

[Example 13]
Experiments were conducted in the same manner as in Examples 1 to 12 in order to observe the hydrogen absorption and release reactions. However, the peripheral speed of the roll made of Cu was set to 10 m / s, and a super-quenched ribbon of Sm (Fe _0.16 Co _0.12 Cu _0.72 ) ₄ having a 1-4 composition as a subphase was produced. This was heat-treated up to 700 ° C. in a hydrogen stream, and the exhaust gas was measured. That is, if there is hydrogen absorption, the exhaust gas decreases, and if there is hydrogen release, the exhaust gas increases.

  FIG. 5 shows the change in the amount of hydrogen in the exhaust gas with temperature. FIG. 5 shows that the subphase absorbs hydrogen in the temperature range from 250 ° C. to 400 ° C. and releases hydrogen in the temperature range from 540 ° C. to 660 ° C. Such absorption and release of hydrogen reflects that the secondary phase is a rare earth-rich phase with a relatively large amount of rare earths.

[Example 14]
<Process A>
The raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of the evaporation of rare earth elements during melting. Specifically, an ultra-quenched ribbon having the same composition as the ultra-quenched ribbon produced in Example 6 was produced under the same quenching conditions.
<Process B>
The ultra-quenched ribbon produced in step A was included in the Nb foil, and heat treatment was performed at 1050 ° C. for 20 minutes in an Ar flow.
<Process C>
Heat treatment was performed up to 500 ° C. in a flowing hydrogen gas to sufficiently absorb hydrogen, and then switched to a vacuum atmosphere to sufficiently release hydrogen and cooled to room temperature. The obtained alloy was confirmed to have cracks between the main phase and the subphase or in the subphase.

  The alloy for rare earth magnets which passed through the process A to the process C and the alloy for rare earth magnets which passed through only the process A and the process B were each grind | pulverized with the jet mill. The resulting fine powder was oriented in a magnetic field and the saturation magnetization was measured to evaluate the degree of orientation. As a result, it was found that the rare earth magnet alloy that had undergone the process A to the process C showed a higher magnetic flux density and a higher degree of orientation than the rare earth magnet alloy that had undergone only the process A and the process B. . This is because the probability of cracking at the main phase crystal grain edge portion is increased mainly by introducing cracks in the main phase and subphase or subphase.

[Examples 15 to 18]
<Process A>
An alloy was prepared using a strip casting apparatus. First, raw materials of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of evaporation of rare earth elements during melting. These raw material metals were put into a silica crucible, and Ar was heated to 1500 ° C. by high frequency induction heating in an atmosphere of −50 kPa at a gauge pressure to dissolve the raw material. Thereafter, the temperature of the molten metal was lowered to 1450 ° C., temporarily stored in a tundish, and then poured onto a copper cooling roll rotating at a peripheral speed of 1.5 m / s to be cooled. The cooled alloy was crushed by a crusher installed at the lower part of the cooling roll. Thus, a strip cast alloy having a composition Y _0.36 Sm _0.64 (Fe _0.83 Co _0.17 ) _9.76 Ti _0.53 Cu _0.53 was produced. This composition was within the scope of the present invention.
<Process B>
The strip cast alloy produced in step A was put in a molybdenum vessel and heat-treated at 1100 ° C. for 1 hour in an Ar flow.
<Process C>
The alloy obtained in step B is held in a hydrogen atmosphere at a temperature T1 (T1 = 150 ° C., 350 ° C., 600 ° C.) for 1 hour, then switched to an Ar atmosphere and held for 10 minutes, and then to room temperature. Cooled down.

  After filling the cross section of the alloy obtained in Step B and Step C with resin, polishing and cross section polishing were performed, and a field of view of 60 μm × 45 μm was photographed with SEM. The results are shown in FIGS. The number of cracks between main phase grains that can be visually confirmed using an image with a field of view of 60 μm × 45 μm was counted.

  The alloy obtained in Step B and Step C was pulverized in a mortar, and powder X-ray diffraction measurement was performed using powder passed through a mesh having an opening of 75 μm. The ratio of the α- (Fe, Co, Ti) phase was determined from the obtained diffraction pattern by Rietveld analysis. DIFFRACplus Professional TOPAS 4 (manufactured by Bruker AEX Co., Ltd.) was used as the Rietveld analysis software. For comparison of peak height, background was removed using analysis software EVA (manufactured by Bruker Ax Co., Ltd.).

  In the alloy of Example 15 without hydrogen treatment, there was only one crack between main phase grains in the 60 μm × 45 μm field of view. Even in the alloy after the hydrogen treatment at 150 ° C. in Example 16, there was only one crack between the main phase grains. In the alloy of Example 17 after the hydrogen treatment at 350 ° C., cracks between main phase grains were observed at 13 places, which was significantly increased from Examples 15 and 16. In the alloy after the hydrogen treatment at 600 ° C. in Example 18, cracks between main phase grains were observed at three places, and a fine structure was observed in the outer shell of the main phase grains. The α- (Fe, Co, Ti) phase ratio of Example 18 is 34.3 wt%, which is very high compared to Examples 15 to 17, and the fine structure of the main phase grain outer shell of Example 18 is mainly It can be said that this is a structure containing an α- (Fe, Co, Ti) phase formed by phase decomposition. As described above, when the α- (Fe, Co, Ti) phase ratio is high, the main phase is sufficiently regenerated by the recombination reaction even if heat treatment is performed in a subsequent vacuum atmosphere or inert gas atmosphere. Therefore, it is desirable that the α- (Fe, Co, Ti) phase ratio is low (preferably 10% or less, more preferably 5% or less).

FIG. 6 is a diagram illustrating powder X-ray diffraction patterns of the samples of Examples 15 to 18 in the embodiment of the present disclosure from the lower stage, respectively. As shown in FIG. 6, in Examples 15 and 16 (in FIG. 6, no hydrogen treatment is Example 15 and T1 = 150 ° C. is Example 16), the lattice spacing value is in the range of 2.50 mm or more and 2.57 mm or less. A clear (121) peak of KHg type ₂ was observed. On the other hand, in Example 17 (in FIG. 6, T1 = 350 ° C.), no clear (121) peak of KHg ₂ type was observed, and the lattice spacing value was 3.03 to 3.10 Å instead. As a result, the peak intensity of the unidentified phase existing in the range of () was higher (position of ■ in FIG. 6). A BSE image with a field of view of 240 μm × 180 μm was taken. Then, using the image analysis software Scandium, the area ratio of the subphase in the image with a field of view of 240 μm × 180 μm is 4.4 in Example 15 before the hydrogen treatment, and after 350 ° C. hydrogen treatment. 17 was 6.4.

  From these results, the lattice spacing value in the powder X-ray diffraction pattern is 3 in order to obtain highly-orientable anisotropic magnetic powder containing a large amount of fine powder of a single crystal unit by increasing the probability that the magnetic powder breaks at the crack portion. The maximum value of the X-ray intensity in the range of 0.03 to 3.10 mm is larger than the maximum value of the X-ray intensity in the range of the lattice spacing value of 2.50 to 2.57; It is preferable that the α- (Fe, Co, Ti) phase is 10% by weight or less (Example 17).

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

Claims (14)

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 composition formula (1):
_{R (Fe, Co) w-} z Ti z Cu α (1)
R is at least one rare earth element,
w, z, and α are respectively 8 ≦ w ≦ 13,
0.42 ≦ z <0.70, and 0.40 ≦ α ≦ 0.70,
Satisfying rare earth magnet alloys. R is composed of R1 and R2,
The overall composition is represented by the following composition formula (2):
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α (2)
R1 is Y or Y and Gd, Y is 50 mol% or more of the entire R1,
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, Sm is necessarily included, and Sm is 50 mol% or more of the entire R2,
x and y are respectively
0.5 ≦ x ≦ 1.0,
0 ≦ y ≦ 0.4,
The alloy for rare earth magnets according to claim 1, wherein The main phase has a ThMn ₁₂ type crystal structure,
The composition of the main phase is represented by the following composition formula (3):
R1 _{1-x ′} R2 _{x ′} (Fe _{1-y ′} Co _{y ′} ) _{12-z′-α ′ Tiz ′} 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
The alloy for rare earth magnets according to claim 1 or 2, satisfying 0.24≤α'≤0.37. z ′ is 0.48 ≦ z ′ <0.74.
The alloy for rare earth magnets according to claim 3, wherein: The main phase is a phase having a ThMn ₁₂ type crystal structure,
The rare earth magnet alloy according to any one of claims 1 to 4, wherein the subphase is mainly a crystal phase having a Cu composition of 50 mol% or more of the entire subphase. The rare earth magnet alloy according to claim 5, wherein the subphase has a crystal structure of KHg ₂ type. The rare earth magnet alloy according to claim 6, wherein the phase of the KHg ₂ type crystal structure is 50% or more of the subphase by volume ratio.   The subphase includes R atoms, and the R atoms present in the subphase have a molar ratio [R2] / ([R1] + [R2]) higher than the composition of the entire alloy. The alloy for rare earth magnets according to claim 1.   The rare earth magnet alloy according to any one of claims 5 to 8, wherein a weight ratio of the subphase in the entire rare earth magnet alloy is 3 wt% or more and 10 wt% or less.   The rare earth magnet alloy according to claim 5, wherein the subphase absorbs and releases hydrogen.   The rare earth magnet alloy according to claim 10, wherein hydrogen absorption and release occur at a temperature of at least 700 ° C. or less.   The rare earth magnet alloy according to any one of claims 1 to 11, wherein a crack is generated between or in the main phase and the sub phase.   In the powder X-ray diffraction pattern, the maximum value of the X-ray intensity in the range of the lattice plane interval value of 3.03 mm to 3.10 mm is in the range of the lattice plane distance value of 2.50 mm to 2.57 mm. The rare earth magnet alloy according to any one of claims 1 to 12, wherein the alloy is greater than the maximum value of the X-ray intensity and the α- (Fe, Co, Ti) phase is 10 wt% or less.   The rare earth magnet alloy according to claim 13, wherein the α- (Fe, Co, Ti) phase is 5 wt% or less.