JP7843935B2 - Rare earth sintered magnets, methods for manufacturing rare earth sintered magnets, rotors and rotating machines - Google Patents

Description

This disclosure relates to a rare earth sintered magnet, which is a permanent magnet obtained by sintering a material containing rare earth elements; a method for manufacturing a rare earth sintered magnet; a rotor; and a rotating machine.

R _- T-B permanent magnets, which have a tetragonal _R₂T₁₄B intermetallic compound as the main phase, are known. Here, R is a rare earth element, T is a transition metal element such as Fe (iron) or Fe partially substituted with Co (cobalt), and B is boron. R-T-B permanent magnets are used in various high-value-added components, including industrial motors. In particular, Nd-Fe-B sintered magnets, where R is Nd (neodymium), are used in various components due to their excellent magnetic properties. Furthermore, since industrial motors are often used in high-temperature environments exceeding 100°C, attempts are being made to improve coercivity by adding heavy rare earth elements such as Dy (dysprosium) to Nd-T-B sintered magnets.

In recent years, the production of Nd-Fe-B sintered magnets has expanded, leading to increased consumption of Nd and heavy rare earth elements such as Dy and Tb (terbium). However, Nd and heavy rare earth elements are expensive and their regional distribution is highly uneven, posing procurement risks. Therefore, technologies to reduce the consumption of Nd and heavy rare earth elements are being researched.

Patent Document 1 discloses an _R -T-B sintered magnet containing main phase particles made of _R2T14B crystal, where R is one or more rare earth elements, with the heavy rare earth element RH being essential; T is one or more transition metal elements, with Fe or Fe and Co being essential; and B is boron. A portion of the main phase particles contains a plurality of low heavy rare earth element crystal phases and a plurality of nonmagnetic R-rich phases internally. Here, the low heavy rare earth element crystal phase consists of _R2T14B crystal, and the concentration of the heavy rare earth element is relatively low compared to the concentration of the heavy rare earth element in the entire main phase particle. The nonmagnetic R-rich phase is a phase in which the R content is 70 at.% or more _and 100 at.% or less. Furthermore, a portion of the main phase particles has a core-shell structure having a core portion and a shell portion surrounding the core portion, in which the total heavy rare earth element concentration is lower than that of the core portion. According to the technology described in Patent Document 1, an R-T-B sintered magnet can be obtained that has improved coercivity and is low cost.

Patent Document 2 discloses a method for manufacturing a _{rare earth} magnet, comprising: a first step of manufacturing a sintered body having a structure consisting of _a main phase and a grain boundary _phase , represented by the compositional formula (R1 _1-x R2 _x ) a TM b B c M d; a second step of manufacturing a rare earth magnet precursor by subjecting the sintered body to hot plastic deformation; and a third step of manufacturing a rare earth magnet by diffusing and permeating the molten R3-M modified alloy into the grain boundary phase of the rare earth magnet precursor. Here, R1 is one or more rare earth elements including Y (yttrium), R2 is a rare earth element different from R1, TM is a transition metal including one or more of Fe, Ni (nickel), and Co, B is boron, and M is one or more of Ti (titanium), Ga (gallium), Zn (zinc), Si (silicon), Al (aluminum), Nb (niobium), Zr (zirconium), Ni, Co, Mn (manganese), V (vanadium), W (tungsten), Ta (tantalum), Ge (germanium), Cu (copper), Cr (chromium), Hf (hafnium), Mo (molybdenum), P (phosphorus), C (carbon), Mg (magnesium), Hg (mercury), Ag (silver), and Au (gold). Furthermore, x, a, b, c, and d are 0.01 ≤ x ≤ 1, 12 ≤ a ≤ 20, b = 100 - a - c - d, 5 ≤ c ≤ 20, 0 ≤ d ≤ 3, and all are at. %. Also, R3 is a rare earth element containing R1 and R2. Hereafter, the hot plastic deformation applied to the sintered body will be referred to as hot working. According to the technology described in Patent Document 2, it is possible to manufacture rare earth magnets that reduce heavy rare earth elements and exhibit excellent coercivity as well as magnetization, even when the main phase fraction is high.

Japanese Patent Publication No. 2018-174313 Japanese Patent Publication No. 2015-153813

However, while the R-T-B sintered magnet described in Patent Document 1 can improve coercivity because a phase containing heavy rare earth elements is present in the main phase, the microstructure of the R-T-B sintered magnet described in Patent Document 1 does not allow for the residual magnetic flux density required for industrial motors, etc., and there is a possibility that the magnetic properties will deteriorate under thermal load. Furthermore, while the rare earth magnet manufactured by the manufacturing method described in Patent Document 2 can reduce the amount of heavy rare earth elements and improve coercivity, the manufacturing method includes hot working, which reduces the grain size of the main phase, resulting in a problem of deteriorated magnetization performance.

This disclosure has been made in view of the above, and aims to provide a rare-earth sintered magnet that can improve coercivity without reducing residual magnetic flux density and magnetization performance compared to conventional magnets.

To solve the above-mentioned problems and achieve the objective, the rare earth sintered magnet according to this disclosure comprises a main phase containing crystal grains based on the Nd 2 Fe 14 B crystal structure, satisfying the general formula (Nd, Pr (praseodymium)), R (R) - Fe - B, when RH is a heavy rare earth element including at least one of Dy and Tb, and R is one or more rare earth elements selected from Nd _, Pr ( _praseodymium ), Dy, and Tb, and a subphase existing between a plurality of main phases. The main phase has a core portion and a shell portion covering the core portion. The main phase has a first main phase where CNd > CPr and a second main phase where CNd < CPr, when the concentration of Nd in the core portion is CNd and the concentration of Pr in the core portion is CPr. The concentration of the heavy rare earth element RH in the core portion of the first main phase is higher than the concentration of the heavy rare earth element RH in the core portion of the second main phase. The first and second main phases are mixed, and heavy rare earth elements are present on at least a portion of the surface of both the first and second main phases.

The rare-earth sintered magnet according to this disclosure has the effect of improving coercivity without reducing the residual magnetic flux density and magnetization performance compared to conventional magnets.

This figure schematically shows an example of the structure of a rare-earth sintered magnet in its sintered state according to Embodiment 1. This figure schematically shows an example of the structure of a rare-earth sintered magnet in its sintered state according to Embodiment 2. This figure schematically shows an example of the structure of a rare-earth sintered magnet in its sintered state according to Embodiment 3. Diagram showing atomic sites in the tetragonal _{Nd₂Fe₁₄B} crystal structure _. A flowchart showing an example of the procedure for manufacturing a rare earth sintered magnet according to Embodiment 4. A flowchart showing an example of the procedure for the manufacturing process of a rare earth sintered magnet alloy according to Embodiment 4. A flowchart showing an example of the procedure for the diffusion precursor manufacturing process according to Embodiment 4. A schematic cross-sectional view showing an example of the configuration of a rotor equipped with a rare-earth sintered magnet according to Embodiment 5. A schematic cross-sectional view showing an example of the configuration of a rotating machine according to Embodiment 6.

The rare earth sintered magnet, the method for manufacturing the rare earth sintered magnet, the rotor, and the rotating machine according to embodiments of this disclosure will be described in detail below with reference to the drawings.

Embodiment 1.
Figure 1 is a schematic diagram showing an example of the structure of a rare earth sintered magnet in its sintered state according to Embodiment 1. The rare earth sintered magnet 1 according to Embodiment 1 has a main phase 10 that satisfies the general formula (Nd, Pr, RH, R)-Fe-B and contains crystal grains based on the Nd ₂ Fe ₁₄ B crystal structure, and the main phase 10 has a core portion and a shell portion that covers the core portion. Here, RH is a heavy rare earth element, and in one example it is Dy, Tb, Gd (gadolinium), Ho (holmium), and preferably RH is a heavy rare earth element containing at least one of Dy and Tb. R is one or more rare earth elements selected from Nd, Pr, and RH. The shell portion has a different composition from the core portion and is provided to cover the core portion. The rare earth sintered magnet 1 also further has a sub-phase 20 that exists between the main phases 10, that is, between a plurality of main phases 10. The secondary phase 20 is a phase based on an oxide phase whose main component is (Nd, Pr, R)-O, where O is oxygen.

In the rare earth sintered magnet 1 according to Embodiment 1, when the concentration of Nd in the core portions 11c and 12c is CNd and the concentration of Pr in the core portions 11c and 12c is CPr, the main phase 10 has a first main phase 11 where CNd > CPr and a second main phase 12 where CNd < CPr, and the first main phase 11 and the second main phase 12 are mixed together. Furthermore, the concentration of heavy rare earth element RH in the core portion 11c of the first main phase 11 is higher than the concentration of heavy rare earth element RH in the core portion 12c of the second main phase 12. That is, if the concentration of heavy rare earth element RH in the core portion 11c of the first main phase 11 is C1RH and the concentration of heavy rare earth element RH in the core portion 12c of the second main phase 12 is C2RH, then C1RH > C2RH. The first main phase 11 has a core portion 11c and a shell portion 11s that has a different composition from the core portion 11c and covers the core portion 11c. The second main phase 12 has a core portion 12c and a shell portion 12s that has a different composition from the core portion 12c and covers the core portion 12c.

Alternatively, in the rare earth sintered magnet 1 according to Embodiment 1, when the sum of the concentration of Nd and the concentration of heavy rare earth element RH in the core portions 11c and 12c is defined as C(Nd,RH), the main phase 10 has a first main phase 11 where C(Nd,RH) > CPr and a second main phase 12 where C(Nd,RH) < CPr, and it can be said that the first main phase 11 and the second main phase 12 are mixed together.

In other words, the rare earth sintered magnet 1 has two types of main phases 10: a first main phase 11 and a second main phase 12. Focusing on the core portions 11c and 12c of the two types of main phases 10, in the first main phase 11, the sum of the concentration of Nd and the concentration of heavy rare earth element RH is higher than the concentration of Pr, while conversely, in the second main phase 12, the concentration of Pr is higher than the sum of the concentration of Nd and the concentration of heavy rare earth element RH. By mixing two types of main phases 10 having core-shell structures with different anisotropic magnetic fields, i.e., different magnetic anisotropy, and selectively including heavy rare earth element RH within the main phases 10, it is possible to reduce Nd and heavy rare earth element RH while maintaining good magnetization and improving residual magnetic flux density and coercivity. Furthermore, since the coercivity is greatly improved by adding heavy rare earth element RH, it also contributes to suppressing the deterioration of magnetic properties due to temperature changes.

The concentration difference shown here, "C1RH > C2RH," indicates that, as determined by mapping analysis using an electron probe microanalyzer (EPMA), there is a clear difference in the detection intensity of heavy rare earth element RH between the core portion 11c of the first main phase 11 and the core portion 12c of the second main phase 12. Specifically, for the concentration of heavy rare earth element RH in the core portion 11c of the first main phase 11, the detection intensity of the EPMA is higher than the average detection intensity of heavy rare earth element RH, while for the concentration of heavy rare earth element RH in the core portion 12c of the second main phase 12, the detection intensity of the EPMA is near the lower limit of the detection intensity of heavy rare earth element RH. In one example, heavy rare earth element RH is present in the core portion 11c of the first main phase 11, but hardly present in the core portion 12c of the second main phase 12.

Furthermore, the concentration difference shown between the first main phase 11, where C(Nd,RH) > CPr, and the second main phase 12, where C(Nd,RH) < CPr, indicates a clear difference between the detection intensity of Nd and heavy rare earth element RH and the detection intensity of Pr, as revealed by mapping analysis using EPMA. Specifically, taking the first main phase 11 as an example, the detection intensity of Nd and heavy rare earth element RH in the core portion 11c is higher than the average detection intensity of Nd and heavy rare earth element RH, while the detection intensity of Pr is near the lower limit of the detection intensity of Pr. The second main phase 12 is the opposite of the first main phase 11.

Furthermore, in the rare earth sintered magnet 1 according to Embodiment 1, when the Nd concentration in the core portion 11c of the first main phase 11 is C1Nd, the Nd concentration in the core portion 12c of the second main phase 12 is C2Nd, the Pr concentration in the core portion 11c of the first main phase 11 is C1Pr, and the Pr concentration in the core portion 12c of the second main phase 12 is C2Pr, the relationship C1Nd > C2Nd and C1Pr < C2Pr is satisfied. In other words, the Nd concentration is higher in the core portion 11c of the first main phase 11 than in the core portion 12c of the second main phase 12, and conversely, the Pr concentration is higher in the core portion 12c of the second main phase 12 than in the core portion 11c of the first main phase 11. This difference in concentration also means that there is a difference in the detection intensity of Nd and Pr in the mapping analysis using EPMA as described above. Specifically, in the case of Nd concentration, this means that the detection intensity of Nd in the EPMA of the core portion 11c of the first main phase 11 is higher than the average detection intensity of Nd, and the detection intensity of Nd in the EPMA of the core portion 12c of the second main phase 12 is lower than the average detection intensity of Nd. In the case of Pr concentration, this means that the detection intensity of Pr in the EPMA of the core portion 12c of the second main phase 12 is higher than the average detection intensity of Pr, and the detection intensity of Pr in the EPMA of the core portion 11c of the first main phase 11 is lower than the average detection intensity of Pr. In other words, there is a lot of Pr in the core portion 12c of the second main phase 12 where the Nd concentration is low, and conversely, there is a lot of Nd in the core portion 11c of the first main phase 11 where the Pr concentration is low. By controlling the microstructure to this extent, it is possible to obtain a rare-earth sintered magnet 1 with excellent magnetic properties.

Furthermore, in the rare earth sintered magnet 1 according to Embodiment 1, there are more first main phases 11 having C(Nd,RH) > CPr than second main phases 12 having C(Nd,RH) < CPr. In other words, the number of first main phases 11 having the composition formula (Nd,RH) ₂ Fe ₁₄ B is greater than the number of second main phases 12 having the composition formula Pr ₂ Fe ₁₄ B. This is because increasing the amount of first main phases 11 having the composition formula (Nd,RH) ₂ Fe ₁₄ B yields better magnetic and thermal properties than increasing the amount of second main phases 12 having the composition formula Pr ₂ Fe ₁₄ B. Moreover, by controlling the microstructure in this way, overall grain refinement is also suppressed, making it possible to obtain superior magnetic properties compared to conventional methods while ensuring magnetization.

Furthermore, in the rare earth sintered magnet 1 according to Embodiment 1, focusing on the shell portions 11s and 12s of the core-shell structure, when the concentration of Nd in the shell portions 11s and 12s is set to SNd, the concentration of Pr in the shell portions 11s and 12s is set to SPr, and the concentration of the heavy rare earth element RH in the shell portions 11s and 12s is set to SRH, the first main phase 11 satisfies the relationships CNd > SNd, CPr < SPr, CRH > SRH, and the second main phase 12 satisfies the relationships CNd < SNd, CPr > SPr, CRH < SRH. Specifically, in the shell portion 11s of the first main phase 11, the concentrations of Nd and the heavy rare earth element RH are lower than in the core portion 11c, but the concentration of Pr is higher than in the core portion 11c. Furthermore, in the shell portion 12s of the second main phase 12, the concentration of Pr is lower than in the core portion 12c, while the concentrations of Nd and heavy rare earth elements RH are higher than in the core portion 12c.

In other words, in the rare earth sintered magnet 1 according to Embodiment 1, focusing on the shell portions 11s and 12s of the core-shell structure, if we define S(Nd,RH) as the sum of the concentrations of Nd and heavy rare earth element RH in the shell portions 11s and 12s, and SPr as the concentration of Pr in the shell portions 11s and 12s, then the first main phase 11 satisfies the relationship C(Nd,RH) > S(Nd,RH) and CPr < SPr, and the second main phase 12 satisfies the relationship C(Nd,RH) < S(Nd,RH) and CPr > SPr. Specifically, in the shell portion 11s of the first main phase 11, the sum of the concentrations of Nd and heavy rare earth element RH is smaller than that of the core portion 11c, while the concentration of Pr is higher than that of the core portion 11c. Furthermore, in the shell portion 12s of the second main phase 12, the concentration of Pr is lower than in the core portion 12c, but the sum of the concentrations of Nd and heavy rare earth element RH is higher than in the core portion 12c.

By forming a main phase 10 having a shell portion 11s with a high concentration of Pr, as in the first main phase 11, the coercivity can be improved. Furthermore, by forming a main phase 10 having a shell portion 12s with a high concentration of Nd and a high sum of heavy rare earth element RH, as in the second main phase 12, the decrease in residual magnetic flux density can be suppressed while maintaining coercivity. By selectively controlling the microstructure to achieve such a configuration, the rare earth sintered magnet 1 can exhibit superior magnetic properties compared to conventional magnets.

Furthermore, the main phase 10 has a heavy rare earth element-containing layer 31 that contains heavy rare earth elements RH in at least a portion of its surface. In other words, heavy rare earth elements RH are present in at least a portion of the surface of the main phase 10, i.e., the first main phase 11 and the second main phase 12. More specifically, heavy rare earth elements RH are present in the core portions 11c and 12c, but even more heavy rare earth elements RH are present in at least a portion of the outer circumferential surface of the shell portions 11s and 12s. In this way, the coercivity is increased by the presence of heavy rare earth elements RH in the R sites of the first main phase 11 and the second main phase 12. To obtain such an effect, it is desirable that the proportion of heavy rare earth elements RH in the main phase 10 be greater than 0 at.% and 10 at.% or less.

Furthermore, the average grain size of the crystal grains in the main phase 10 is preferably 100 μm or less, and more preferably 0.5 μm to 50 μm to improve magnetic properties. Moreover, by setting it to approximately 1 μm to 10 μm, the grain size differs from that of the microstructure produced by hot working, maintaining good magnetization performance and making it possible to produce a rare-earth sintered magnet 1 with superior magnetic properties compared to conventional magnets.

The rare earth sintered magnet 1 according to Embodiment 1 may contain an additive element M that further improves the magnetic properties. The additive element M is one or more elements selected from the group Ga, Cu, Al, Co, Zr, Ti, Nb, and Mn. Therefore, if RH is a heavy rare earth element selected from the group Dy, Tb, Gd, and Ho, and R is Nd, Pr, and a rare earth element other than the heavy rare earth element RH, the general formula of the rare earth sintered magnet 1 according to Embodiment 1 can be expressed as (Nd _a Pr _b R _c RH _d )Fe _e B _f M _g . It is desirable that a, b, c, d, e, f, and g satisfy the following relationship.

5 ≤ a + b ≤ 20
0<c+d<(a+b)
0 < d < 10
70 ≤ e ≤ 90
0.5 ≤ f ≤ 10
0 ≤ g ≤ 5
a+b+c+d+e+f+g=100at. %

The rare earth sintered magnet 1 according to Embodiment 1 satisfies the general formula (Nd, Pr, RH, R)-Fe-B when R is one or more rare earth elements selected from those other than Nd, Pr and the heavy rare earth element RH, and the main phase 10, which contains crystal grains based on the Nd ₂ Fe ₁₄ B crystal structure, has core portions 11c, 12c and shell portions 11s, 12s covering the core portions 11c, 12c, and the main phase 10 has a first main phase 11 where CNd > CPr and a second main phase 12 where CNd < CPr, and the concentration of the heavy rare earth element RH in the first main phase 11 is higher than the concentration of the heavy rare earth element RH in the second main phase 12, and the first main phase 11 and the second main phase 12 are mixed. This configuration makes it possible to obtain a rare-earth sintered magnet 1 with improved magnetic properties and magnetization compared to conventional magnets, while reducing the use of Nd and heavy rare-earth elements RH.

This document compares the R-T-B sintered magnet described in Patent Document 1 with the rare-earth sintered magnet 1 according to Embodiment 1. The R-T-B sintered magnet described in Patent Document 1 has one or more rare-earth elements, with heavy rare-earth element RH being essential, and has one type of main phase particle composed of a core portion and a shell portion. In other words, in Patent Document 1, heavy rare-earth element RH is contained in all main phase particles of the R-T-B sintered magnet. On the other hand, in the rare-earth sintered magnet 1 according to Embodiment 1, the main phase 10 has a first main phase 11 and a second main phase 12, and the concentration of heavy rare-earth element RH is lower in the second main phase 12 compared to the first main phase 11. In one example, the core portion 11c of the first main phase 11 contains heavy rare-earth element RH, but the core portion 12c of the second main phase 12 contains almost no heavy rare-earth element RH. In other words, heavy rare-earth element RH is selectively arranged within the main phase 10. Thus, compared to the R-T-B sintered magnet described in Patent Document 1, in which it is necessary to ensure that all of the single main phase contains heavy rare earth elements RH, the rare earth sintered magnet 1 according to Embodiment 1, which has a main phase 10 comprising a first main phase 11 and a second main phase 12 having a lower concentration of heavy rare earth elements RH compared to the first main phase 11, can reduce the amount of heavy rare earth elements RH used. In particular, when the second main phase 12 contains almost no heavy rare earth elements RH, the heavy rare earth elements RH are selectively arranged, and the amount of heavy rare earth elements RH used can be reduced compared to the R-T-B sintered magnet described in Patent Document 1.

This document compares a rare earth magnet manufactured using the technology described in Patent Document 2 with a rare earth sintered magnet 1 according to Embodiment 1. To obtain magnetic properties equivalent to those of the rare earth sintered magnet 1 according to Embodiment 1 using a rare earth magnet manufactured using the technology described in Patent Document 2, a larger amount of heavy rare earth element RH must be added compared to the rare earth sintered magnet 1 according to Embodiment 1, as shown in the examples described later. In other words, the rare earth sintered magnet 1 according to Embodiment 1 can use less heavy rare earth element RH compared to the technology described in Patent Document 2 when attempting to obtain the same magnetic properties. Alternatively, if the heavy rare earth element RH content in the rare earth magnet manufactured using the technology described in Patent Document 2 is the same as the heavy rare earth element RH content in the rare earth sintered magnet 1 according to Embodiments 1 and 2, the magnetic properties of the rare earth magnet manufactured using the technology described in Patent Document 2 will be lower than those of the rare earth sintered magnet 1 according to Embodiments 1 and 2.

In the rare earth sintered magnet 1 according to Embodiment 1, the first main phase 11 and the second main phase 12 are configured to satisfy the relationships C1Nd > C2Nd and C1Pr < C2Pr. Alternatively, the number of first main phases 11 is greater than the number of second main phases 12. Alternatively, the first main phase 11 satisfies the relationships CNd > SNd, CPr < SPr, CRH > SRH, and the second main phase 12 satisfies the relationships CNd < SNd, CPr > SPr, CRH < SRH. This also makes it possible to obtain a rare earth sintered magnet 1 with improved magnetic properties and magnetization while suppressing the use of Nd and heavy rare earth elements RH.

Furthermore, heavy rare earth elements RH are present on at least a portion of the surfaces of the first main phase 11 and the second main phase 12. That is, the first main phase 11 and the second main phase 12 each have a heavy rare earth element-containing layer 31 on at least a portion of their respective surfaces. This also allows for the production of a rare earth sintered magnet 1 that improves coercivity compared to conventional magnets while suppressing a significant decrease in residual magnetic flux density, all while reducing the use of heavy rare earth elements RH. In other words, it has the effect of improving the magnetic properties of the rare earth sintered magnet 1 compared to conventional magnets.

Furthermore, by including the heavy rare earth element RH in the main phase 10, a rare earth sintered magnet 1 with significantly improved coercivity compared to conventional magnets can be obtained. Also, as shown in the examples described later, a rare earth sintered magnet 1 with a better temperature coefficient of coercivity compared to conventional magnets can be obtained. Therefore, even when a thermal load is applied to the rare earth sintered magnet 1, the coercivity is greater than in conventional magnets, and the decrease in coercivity due to temperature rise is also more gradual than in conventional magnets. In other words, because the coercivity is significantly improved compared to conventional rare earth sintered magnets, the magnetic properties of the rare earth sintered magnet 1 when subjected to a thermal load are also better than in conventional magnets.

Embodiment 2.
Figure 2 is a schematic diagram showing an example of the structure of a rare earth sintered magnet in its sintered state according to Embodiment 2. Components identical to those in Embodiment 1 are denoted by the same reference numerals, and their descriptions are omitted. The rare earth sintered magnet 1 according to Embodiment 2 has a main phase 10 and a sub-phase 20.

The main phase 10 has the same structure as in Embodiment 1. That is, the main phase 10 has a first main phase 11 and a second main phase 12 having a core-shell structure, and the composition of the core portions 11c, 12c and the composition of the shell portions 11s, 12s are the same as those described in Embodiment 1. However, in Embodiment 2, the heavy rare earth element-containing layer 31 may or may not be present on the surface of the main phase 10.

The subphase 20 is a phase based on an oxide phase whose main component is (Nd, Pr, RH, R)-O. In other words, in Embodiment 2, the subphase 20 contains the heavy rare earth element RH. The heavy rare earth element RH is distributed throughout the subphase 20. In one example, the heavy rare earth element RH is uniformly distributed within the subphase 20.

Thus, in Embodiment 2, a subphase 20 containing heavy rare earth elements RH exists between the main phase 10 and the main phase 10. The heavy rare earth elements RH are uniformly distributed within the subphase 20, and it can also be considered that the heavy rare earth elements RH are embedded in a part of the surface of the main phase 10 that is in contact with the subphase 20. In other words, it is considered that the heavy rare earth elements RH are embedded in a part of the shell portions 11s and 12s. Therefore, similar to Embodiment 1, it is possible to suppress the decrease in residual magnetic flux density while improving the coercivity of the rare earth sintered magnet 1.

As shown in Figure 2, in the rare earth sintered magnet 1, the main phase 10 is in contact with other main phases 10 without the subphase 20, or in contact with other main phases 10 via the subphase 20. That is, at least a portion of the surface of the main phase 10 is in contact with the subphase 20. Furthermore, the subphase 20 contains the heavy rare earth element RH. Therefore, at least a portion of the surface of the main phase 10 is covered by the subphase 20 containing the heavy rare earth element RH. From this, it can be said that the heavy rare earth element RH is present on at least a portion of the surface of the main phase 10, specifically on the surface of the main phase 10 that is in contact with the subphase 20. In other words, Embodiment 1 shows the distribution of the heavy rare earth element RH focusing on the interface between the main phase 10 and the subphase 20 in the same rare earth sintered magnet 1, while Embodiment 2 shows the distribution of the heavy rare earth element RH focusing on the subphase 20. Thus, Embodiment 1 and Embodiment 2 can be described as the same rare-earth sintered magnet 1 viewed from different angles.

In Embodiment 2, as in Embodiment 1, it is possible to obtain a rare-earth sintered magnet 1 that improves coercivity compared to conventional magnets while suppressing a significant decrease in residual magnetic flux density, while reducing the use of heavy rare-earth elements RH. Alternatively, it is possible to obtain a rare-earth sintered magnet 1 with magnetic properties equivalent to or better than conventional magnets, using a smaller amount of heavy rare-earth elements RH compared to conventional magnets. Furthermore, because the coercivity is greatly improved compared to conventional rare-earth sintered magnets, the magnetic properties when the rare-earth sintered magnet 1 is subjected to a thermal load are also better than conventional magnets. In other words, it has the effect of improving the magnetic properties of the rare-earth sintered magnet 1 compared to conventional magnets.

Embodiment 3.
Figure 3 is a schematic diagram showing an example of the structure of the sintered state of a rare earth sintered magnet according to Embodiment 3. The rare earth sintered magnet 1 according to Embodiment 3 has a main phase 10 and a sub-phase 20. The main phase 10 includes a first main phase 11 and a second main phase 12, as described in Embodiment 1. The sub-phase 20 is located between the main phases 10.

Embodiment 3 of the rare earth sintered magnet 1 shows the case where La (lanthanum) and Sm (samarium) are selected as the rare earth element R. When La and Sm are selected as the rare earth element R, the effect of improving magnetic properties and having superior magnetization compared to conventional magnets is further enhanced while suppressing the use of Nd and the heavy rare earth element RH. In this example, the main phase 10 has the composition formula (Nd, Pr, RH, La, Sm) ₂ Fe ₁₄ B. The reason why the rare earth element R of the rare earth sintered magnet 1 having a tetragonal R ₂ Fe ₁₄ B crystal structure is to include rare earth elements La and Sm is that, from the calculation results of magnetic interaction energy using molecular orbital method, a practical rare earth sintered magnet 1 can be obtained that can greatly suppress the decrease in magnetic properties with increasing temperature by adding La and Sm. Furthermore, by intentionally segregating La and Sm at grain boundaries, which are an example of the subphase 20, Nd and Pr are relatively diffused into the main phase 10, thereby increasing the crystalline magnetic anisotropy of the main phase 10. As a result, a core-shell structure is formed in the main phase 10 in which parts with high magnetic anisotropy and parts with low magnetic anisotropy exist, and a first main phase 11 in which CNd > CPr and a second main phase 12 in which CNd < CPr coexist, creating a condition in which a rare earth sintered magnet 1 is easily formed in which the concentration of heavy rare earth elements RH in the core portion 11c of the first main phase 11 is higher than the concentration of heavy rare earth elements RH in the core portion 12c of the second main phase 12.

Furthermore, if the amounts of La and Sm added are too high, the amounts of Nd and Pr, which are elements with high magnetic anisotropy constants and saturation magnetic polarization, will decrease, leading to a decline in magnetic properties. For this reason, when the composition ratios of Nd, Pr, La, and Sm are A, B, C, and D, respectively, it is preferable that (A + B) > (C + D).

In the rare earth sintered magnet 1 according to Embodiment 3, when R = La, Sm, in addition to the first main phase 11 and second main phase 12 in Embodiments 1 and 2, there is a subphase 20. The subphase 20 has a crystalline first subphase 21 based on an oxide phase whose main component is represented as (Nd, Pr, RH, La, Sm)-O, and a crystalline second subphase 22 whose main component is represented as (Nd, Pr, RH, La)-O. The concentration of Sm in the subphase 20 is higher in the first subphase 21 than in the second subphase 22. In other words, the first subphase 21 forms an Sm-enriched portion 41 with a higher Sm concentration than the second subphase 22. This has the effect of suppressing the decrease in magnetic properties not only at room temperature but also with increasing temperature.

Here, "the concentration of Sm is higher in the first subphase 21 compared to the second subphase 22" means that, on average, the detection intensity of Sm is higher in the first subphase 21 than in the second subphase 22, as determined by mapping analysis using EPMA.

The crystalline subphase 20 is a collective term for the crystalline first subphase 21 and the crystalline second subphase 22, and exists between the main phase 10. The crystalline first subphase 21 is represented as (Nd, Pr, RH, La, Sm)-O, and the crystalline second subphase 22 is represented as (Nd, Pr, RH, La)-O. Here, (Nd, Pr, RH, La, Sm) means that some of Nd and Pr are substituted by the heavy rare earth elements RH, La, and Sm. Note that the main component elements are listed in parentheses here, so the first subphase 21 and the second subphase 22 may contain trace amounts of other components in addition to the elements shown in parentheses. For example, the second subphase 22 represented as (Nd, Pr, RH, La)-O contains a trace amount of Sm.

In the rare earth sintered magnet 1 according to Embodiment 3, there is a difference in the concentrations of La and Sm between the main phase 10 and the sub-phase 20, with La and Sm being more segregated in the sub-phase 20 than in the main phase 10. That is, the sum of the concentrations of La in the first sub-phase 21 and the second sub-phase 22 is greater than or equal to the concentration of La in the main phase 10, and the sum of the concentrations of Sm in the first sub-phase 21 and the second sub-phase 22 is greater than or equal to the concentration of Sm in the main phase 10. Specifically, the concentrations of La and Sm in the sub-phase 20 are greater than or equal to the concentrations of La and Sm in the main phase 10. Here, the concentration of La in the main phase 10 is the sum of the concentration of La in the first main phase 11 and the concentration of La in the second main phase 12. In other words, the sum of the concentrations of La in the first subphase 21 and the second subphase 22 is higher than the sum of the concentrations of La in the first main phase 11 and the second main phase 12. Also, the concentration of Sm in the main phase 10 is the sum of the concentration of Sm in the first main phase 11 and the concentration of Sm in the second main phase 12. In other words, the sum of the concentrations of Sm in the first subphase 21 and the second subphase 22 is higher than the sum of the concentrations of Sm in the first main phase 11 and the second main phase 12.

Here, when the La concentration in the main phase 10 is X, the La concentration in the first sub-phase 21 is _X1 , the La concentration in the second sub-phase 22 is _X2 , the Sm concentration in the main phase 10 is Y, the Sm concentration in the first sub-phase 21 is _Y1 , and the Sm concentration in the second sub-phase 22 is _Y2 , the following relationship (1) is satisfied.

1<(Y ₁ +Y ₂ )/Y<(X ₁ +X ₂ )/X (1)

Furthermore, from the viewpoint of improving magnetic properties, the following relationships (2) and (3) are satisfied with respect to the concentrations of Nd and Pr contained in the main phase 10.

(CNd+SNd)>(X+Y)...(2)
(CPr+SPr)>(X+Y)...(3)

In the above, the concentration of La in the main phase 10 was assumed to be the sum of the concentrations of La in the first main phase 11 and the second main phase 12, and the concentration of Sm in the main phase 10 was assumed to be the sum of the concentrations of Sm in the first main phase 11 and the second main phase 12. This indicates that both La and Sm are segregated more towards the subphase 20 than towards the main phase 10. However, when viewed locally, the sum of the concentrations of La and Sm in the first main phase 11 and the second main phase 12, and the sum of the concentrations of La and Sm in the first subphase 21 and the second subphase 22, may not satisfy the above relationship. Therefore, more specifically, the concentration of La in the main phase 10 represents the average of the concentrations of La in the first main phase 11 and the second main phase 12, and the concentration of Sm in the main phase 10 represents the average of the concentrations of Sm in the first main phase 11 and the second main phase 12. In this case, the concentration of La in subphase 20, i.e., the sum of the concentrations of La in the first subphase 21 and the second subphase 22, represents the average of the concentrations of La in the first subphase 21 and the second subphase 22, and the concentration of Sm in subphase 20, i.e., the sum of the concentrations of Sm in the first subphase 21 and the second subphase 22, represents the average of the concentrations of Sm in the first subphase 21 and the second subphase 22.

La, due to its high concentration at grain boundaries during the manufacturing process, particularly the heat treatment process, relatively diffuses Nd and Pr into the main phase 10. As a result, in the rare-earth sintered magnet 1 of Embodiment 3, the Nd and Pr in the main phase 10 are not consumed at grain boundaries, and the crystalline magnetic anisotropy is improved. Similarly, Sm, being present in a high concentration in the sub-phase 20, particularly the first sub-phase 21, compared to the main phase 10, also relatively diffuses Nd into the main phase 10, improving the crystalline magnetic anisotropy, just like La.

As described in Embodiment 2, since the subphase 20 contains heavy rare earth elements RH, the first subphase 21 and the second subphase 22 also contain heavy rare earth elements RH. However, in Embodiment 3, the distribution of heavy rare earth elements RH differs between the first subphase 21 and the second subphase 22. In the second subphase 22, where the Sm concentration is lower than in the first subphase 21, the heavy rare earth elements RH are uniformly distributed within the second subphase 22. On the other hand, in the first subphase 21, which forms the Sm-enriched region 41, the heavy rare earth elements RH are not uniformly distributed within the first subphase 21, but rather selectively distributed between the outer periphery of the first subphase 21 and the Sm-enriched region 41, that is, in the inner periphery of the outer periphery of the first subphase 21. Specifically, the heavy rare earth elements RH are present so as to selectively surround the outer periphery of the Sm-enriched region 41, where the Sm concentration of the first subphase 21 is high. From this, it can be said that the first subphase 21 has an Sm-enriched portion 41 and a heavy rare earth element-containing portion 42 in which heavy rare earth elements RH exist and selectively surround the outer edge of the Sm-enriched portion 41. The outer edge of the first subphase 21 is the boundary between the first subphase 21 and the main phase 10.

Similar to Embodiment 2, a first sub-phase 21 and a second sub-phase 22 containing heavy rare earth elements RH exist between the main phase 10. Therefore, it can be considered that heavy rare earth elements RH are incorporated into a portion of the surface of the main phase 10 that is in contact with the first sub-phase 21 and the second sub-phase 22 containing heavy rare earth elements RH. In other words, it is considered that the heavy rare earth elements RH of the sub-phase 20 are incorporated into a portion of the shell portions 11s and 12s. Therefore, in the rare earth sintered magnet 1 according to Embodiment 3, similar to Embodiment 1, it is possible to suppress the decrease in residual magnetic flux density while improving the coercivity of the rare earth sintered magnet 1.

Furthermore, when the cross-section of the rare earth sintered magnet according to Embodiment 3 is analyzed with a Field Emission-Electron Probe Micro Analyzer (FE-EPMA), it is confirmed that, as shown in Figure 3, there is a main phase 10 having a first main phase 11 in which heavy rare earth element RH is distributed in the core portion 11c, and a second main phase 12 in which heavy rare earth element RH is hardly distributed in the core portion 12c, and that a subphase 20 exists between the main phase 10 and the main phase 10. It is also confirmed that the subphase 20 has a first subphase 21 having an Sm-enriched portion 41, and a second subphase 22 with a lower Sm concentration than the first subphase 21. In the second subphase 22, as described above, Tb, which is a heavy rare earth element RH, is uniformly distributed, while in the first subphase 21, it is confirmed that there is a bias in the distribution of Tb. Furthermore, it is confirmed that heavy rare earth element-containing regions 42 exist so as to selectively surround the Sm-enriched regions 41 in the first subphase 21 where the Sm concentration is high. In addition, it is confirmed that there are almost no heavy rare earth element RHs in the Sm-enriched regions 41, and that the concentration of heavy rare earth element RHs selectively distributed around the Sm-enriched regions 41 is higher than the concentration of heavy rare earth element RHs distributed throughout the second subphase 22.

Next, we will explain at which atomic sites in the _{tetragonal R₂Fe₁₄B} crystal structure La and Sm are substituted. Figure 4 shows the atomic sites in the _{tetragonal Nd₂Fe₁₄B} crystal structure. Note that the crystal structure shown in Figure 4 is, in one example, described in FIG. 1 of Reference 1 listed below. The substituted sites are determined by calculating the stabilization energy due to substitution using band calculations and the molecular field approximation of the Heisenberg model, and then determining the site based on the value of that energy.
(Reference 1) _JFHerbst et al. “Relationships between crystal structure and magnetic properties in _{Nd₂Fe₁₄B} ”. PHYSICAL REVIEW B. 1984, Vol.29, No.7, pp. 4176-4178.

First, we will explain how to calculate the stabilization energy for La. The stabilization energy for La can be determined using the Nd ₈ Fe ₅₆ B ₄ crystal cell by calculating the energy difference between (Nd ₇ La ₁ ) Fe ₅₆ B ₄ + Nd and Nd ₈ (Fe ₅₅ La ₁ ) B ₄ + Fe. The smaller the energy value, the more stable the site is when an atom is substituted there. In other words, La is most likely to be substituted at the atomic site with the lowest energy among the atomic sites. In this calculation, it is assumed that the lattice constant in the tetragonal R ₂ Fe ₁₄ B crystal structure does not change due to differences in atomic radius when La is substituted for the original atom. Table 1 shows the stabilization energy of La at each substitution site when the ambient temperature is changed.

According to Table 1, the stable substitution site for La is the Nd(f) site at temperatures above 1000K, and the Fe(c) site at temperatures of 293K and 500K. The rare earth sintered magnet 1 according to Embodiment 3 is heated to a temperature of 1000K or higher to melt the raw material for the rare earth sintered magnet 1, as will be described later, and then rapidly cooled. For this reason, the raw material for the rare earth sintered magnet 1 is considered to be maintained at a temperature of 1000K or higher, i.e., 727°C or higher, preferably around 1300K, i.e., 1027°C. At this time, it is considered that La is substituted to either the Nd(f) site or the Nd(g) site. Here, it is considered that La is preferentially substituted to the energetically stable Nd(f) site, but substitution to the Nd(g) site, which has a smaller energy difference among the substitution sites for La, is also possible. For this reason, the Nd(g) site is also listed as a candidate substitution site for La.

Furthermore, when the rare earth sintered magnet 1 is manufactured by the manufacturing method described later, although the temperature during sintering is above 1000K, the Fe(c) sites listed in Table 1 are repeatedly maintained in an energetically stable temperature range through the first aging process, second aging process, third aging process, fourth aging process, and cooling process described later. In other words, the substitution of La at the Nd sites of the main phase 10 is maintained in an unstable energy state. That is, in the raw material stage of the rare earth sintered magnet 1, La was mainly substituted at the Nd sites of the main phase 10, but in the manufacturing method described later, by deliberately holding the rare earth sintered magnet 1 repeatedly in an unstable energy state temperature range relative to the Nd sites of the main phase 10, a certain amount of La is selectively released from the Nd sites of the main phase 10, resulting in the segregation of La into the subphase 20. As a result, the main phase 10 promotes the formation of a characteristic structure called a core-shell structure.

Next, we will explain how to calculate the stabilization energy for Sm. The stabilization energy of Sm can be determined by the energy difference between (Nd ₇ Sm ₁ ) Fe ₅₆ B ₄ + Nd and Nd ₈ (Fe ₅₅ Sm ₁ ) B ₄ + Fe. The assumption that the lattice constant in the tetragonal R ₂ Fe ₁₄ B crystal structure does not change due to the substitution of atoms is the same as in the case of La. Table 2 shows the stabilization energy of Sm at each substitution site when the ambient temperature is changed.

According to Table 2, unlike La, the stable substitution site for Sm is the Nd(g) site at all temperatures. While substitution to the energetically stable Nd(g) site is considered to be preferential in Sm, substitution to the Nd(f) site, which has a smaller energy difference, is also possible among the substitution sites for Sm.

When the rare earth sintered magnet 1 is manufactured by the manufacturing method described later, the substitution of the Nd(g) site in the main phase 10 is the most energetically stable. However, as described above, by maintaining the magnet in a temperature range where the substitution of the Nd site in the main phase 10 with La becomes unstable, some Sm is also released from the Nd site in the main phase 10 together with La and segregates into the subphase 20. As a result, there is a concentration difference between the main phase 10 and the subphase 20 in terms of La and Sm concentrations. The sum of the concentrations of La in the first subphase 21 and the second subphase 22 is greater than or equal to the concentration of La in the main phase 10, and the sum of the concentrations of Sm in the first subphase 21 and the second subphase 22 is greater than or equal to the concentration of Sm in the main phase 10. More specifically, the average concentration of La in the first subphase 21 and the second subphase 22 is greater than or equal to the average concentration of La in the first main phase 11 and the second main phase 12, and the average concentration of Sm in the first subphase 21 and the second subphase 22 is greater than or equal to the average concentration of Sm in the first main phase 11 and the second main phase 12. In other words, La and Sm are segregated into the subphase 20.

Comparing La and Sm, from an energetic standpoint, it is clear that La, which is held in a temperature range with an unstable energy state, is overwhelmingly more likely to segregate into the subphase 20. Therefore, in the case of a rare-earth sintered magnet 1 prepared with similar concentrations of La and Sm, the segregation rate of La into the subphase 20 is greater than that of Sm present in the rare-earth sintered magnet 1. Repeated holding in this temperature range creates a concentration difference of Sm, which has a lower segregation rate, in the subphase 20, leading to the formation of the first subphase 21 and the second subphase 22. This promotes the formation of the core-shell structure in the main phase 10.

In this explanation, we will primarily focus on Nd, as shown in Figure 4. However, since Nd and Pr are produced as mixtures, as exemplified by Di (didymium), their energy levels are considered to be close. Therefore, the same can be said when Nd is replaced with Pr. By having both Nd and Pr present, it is possible to form a main phase 10 having two types of core-shell structures.

As described above, the rare earth sintered magnet 1 of Embodiment 3 satisfies the general formula (Nd, Pr, RH, R)-Fe-B when R is one or more rare earth elements selected from other than Nd, Pr, and RH, and has a main phase 10 containing crystal grains based on the Nd ₂ Fe ₁₄ B crystal structure, and the main phase 10 has core portions 11c, 12c and shell portions 11s, 12s covering the core portions 11c, 12c, and when R = La, Sm, it has a sub-phase 20 in addition to the first main phase 11 and second main phase 12 of Embodiment 1. The subphase 20 has a crystalline first subphase 21 based on an oxide phase whose main component is (Nd, Pr, RH, La, Sm)-O, and a crystalline second subphase 22 whose main component is (Nd, Pr, RH, La)-O. The concentration of Sm is higher in the first subphase 21 than in the second subphase 22, and the first subphase 21 has an Sm-enriched region 41 where Sm is selectively distributed. In other words, there are two types of main phase 10 and two types of subphase 20. This makes it possible to provide a rare earth sintered magnet 1 with superior magnetic properties, such as the temperature characteristics of the magnetic properties, compared to conventional magnets. Furthermore, by setting R to La and Sm, the main phase 10 becomes a state in which a first main phase 11 where C(Nd, RH) > CPr and a second main phase 12 where C(Nd, RH) < CPr coexist. In other words, the rare earth sintered magnet 1 has a main phase 10 having two types of first main phases 11 and second main phases 12. Focusing on the core portions 11c and 12c of the two types of main phases 10, it is easy to create a main phase 10 having two types of core-shell structures, where the Nd concentration is higher than the Pr concentration in the first main phase 11, and conversely, the Pr concentration is higher than the Nd concentration in the second main phase 12, and the concentration of heavy rare earth elements RH in the first main phase 11 is higher than the concentration of heavy rare earth elements RH in the second main phase 12. As a result, it is possible to further enhance the effect of improving magnetic properties and having superior magnetization compared to conventional methods while suppressing the use of Nd and heavy rare earth elements RH compared to conventional methods.

Furthermore, in Embodiment 3, similar to Embodiment 1, it is possible to obtain a rare-earth sintered magnet 1 that improves coercivity compared to conventional magnets while suppressing a significant decrease in residual magnetic flux density, while minimizing the use of heavy rare-earth elements RH. In addition, because the coercivity is greatly improved compared to conventional rare-earth sintered magnets, the magnetic properties when the rare-earth sintered magnet 1 is subjected to a thermal load are also better than conventional magnets. In other words, it has the effect of improving the magnetic properties of the rare-earth sintered magnet 1 compared to conventional magnets.

Embodiment 4.
Embodiment 4 describes a method for manufacturing the rare earth sintered magnet 1 described in Embodiments 1, 2, and 3. Figure 5 is a flowchart showing an example of the procedure for manufacturing the rare earth sintered magnet according to Embodiment 4. As shown in Figure 5, the method for manufacturing the rare earth sintered magnet 1 includes a rare earth sintered magnet alloy manufacturing step (step S10) for manufacturing a rare earth sintered magnet alloy that will be used as a raw material for a diffusion precursor, which is a sintered body before the heavy rare earth element RH is diffused into the rare earth sintered magnet 1 containing the heavy rare earth element RH; a diffusion precursor manufacturing step (step S20) for forming the diffusion precursor; a grain boundary diffusion step (step S30) for diffusing the heavy rare earth element RH into the diffusion precursor; and a cooling step (step S40) for cooling the diffusion precursor in which the heavy rare earth element RH has been diffused to obtain the rare earth sintered magnet 1.

First, the details of the rare earth sintered magnet alloy manufacturing process in step S10 will be explained. Figure 6 is a flowchart showing an example of the procedure for the rare earth sintered magnet alloy manufacturing process according to Embodiment 4. First, as shown in Figure 6, the manufacturing process for the rare earth sintered magnet alloy, which is the raw material for the diffusion precursor, includes a melting step (step S11) in which the raw material for the rare earth sintered magnet alloy, which contains the elements constituting the diffusion precursor, is heated to a temperature of 1000K or higher to melt it; a first cooling step (step S12) in which the molten raw material is cooled on a rotating body to obtain a solidified alloy; and a second cooling step (step S13) in which the solidified alloy is further cooled in a container. This makes it possible to manufacture a rare earth sintered magnet alloy. Each step will be explained below.

In the melting step S11, the raw materials for the diffusion precursor are heated to a temperature of 1000K or higher in a crucible in an atmosphere containing an inert gas such as Ar (argon) or in a vacuum to melt them. This prepares a molten alloy in which the rare earth sintered magnet alloy is molten. When manufacturing the rare earth sintered magnet 1 of Embodiments 1 and 2, Nd, Pr, RH, R, Fe, and B can be used as raw materials. When manufacturing the rare earth sintered magnet 1 of Embodiment 3, Nd, Pr, RH, La, Sm, Fe, and B can be used as raw materials. In the case of Embodiment 3, the rare earth element R is La and Sm as in Embodiments 1 and 2. Examples of RH include Dy and Tb. Also, FeB may be used instead of B as a raw material. In this case, one or more elements selected from the group of Ga, Cu, Al, Co, Zr, Ti, Nb, and Mn may be included as additive element M in the raw materials.

Next, in the first cooling step of step S12, the molten alloy prepared in the melting step is poured into a tundish and then onto a single roll, which is a rotating body. As a result, the molten alloy is rapidly cooled on the single roll rotating in a predetermined direction, and a solidified alloy thinner than the ingot alloy is prepared on the single roll from the molten alloy. Here, a single roll is used as the rotating body, but it is not limited to this, and rapid cooling may be achieved by contacting a twin roll, a rotating disk, a rotating cylindrical mold, etc. From the viewpoint of efficiently obtaining a thin solidified alloy, the cooling rate in the first cooling step is preferably 10°C/second or more and ^10⁷ °C/second or less, and more preferably ^10³ °C/second or more and ^10⁴ °C/second or less. The thickness of the solidified alloy is in the range of 0.03 mm or more and 10 mm or less. Solidification of the molten alloy begins from the part that comes into contact with the single roll, and crystals grow in a columnar or needle-like shape in the thickness direction from the contact surface with the single roll. The first cooling step in step S12 corresponds to the first alloy cooling step.

Subsequently, in the second cooling step of step S13, the thin solidified alloy prepared in the first cooling step is placed in a tray container and cooled. The thin solidified alloy breaks down into flake-like rare earth sintered magnet alloy as it enters the tray container and cools. Depending on the cooling rate, ribbon-like rare earth sintered magnet alloy may also be obtained, and it is not limited to flake-like. From the viewpoint of obtaining a rare earth sintered magnet alloy having a microstructure with good temperature characteristics for its magnetic properties, the cooling rate in the second cooling step is preferably ^10⁻² °C/second or more and ^10⁵ °C/second or less, and more preferably ^10⁻¹ °C/second or more and ^10⁻² °C/second or less. The second cooling step of step S13 corresponds to the second alloy cooling step.

The rare earth sintered magnet alloy obtained through these processes has a short-axis size of 3 μm to 10 μm and a long-axis size of 10 μm to 300 μm. In Embodiment 3, it has a fine crystalline structure containing a (Nd, Pr, RH, La, Sm)-Fe-B crystalline phase and a crystalline subphase 20 of an oxide represented by (Nd, Pr, RH, La, Sm)-O. Hereinafter, the crystalline subphase 20 of an oxide represented by (Nd, Pr, RH, La, Sm)-O will be referred to as the (Nd, Pr, RH, La, Sm)-O phase. The (Nd, Pr, RH, La, Sm)-O phase is a non-magnetic phase consisting of an oxide with a relatively high concentration of rare earth elements. The thickness of the (Nd, Pr, RH, La, Sm)-O phase corresponds to the width of the grain boundaries and is 10 μm or less. Because the rare earth sintered magnet alloy produced by the above manufacturing process undergoes a rapid cooling process, its microstructure is finer compared to rare earth sintered magnet alloy obtained by the mold casting method.

Next, the diffusion precursor manufacturing process in step S20 of Figure 5 will be described. Figure 7 is a flowchart showing an example of the procedure for the diffusion precursor manufacturing process according to Embodiment 4. In the following description, the case of manufacturing the rare earth sintered magnet 1 of Embodiment 3 will be used as an example, but the rare earth sintered magnet 1 of Embodiments 1 and 2 can be manufactured by changing the raw materials of the rare earth sintered magnet alloy used. In other words, La and Sm in Embodiment 3 can be Nd, Pr, and rare earth elements R other than the heavy rare earth element RH. As shown in Figure 7, the diffusion precursor manufacturing process includes a grinding step (step S21) in which a rare earth sintered magnet alloy having a (Nd,Pr,RH,La,Sm)-Fe-B crystalline phase and a (Nd,Pr,RH,La,Sm)-O phase is ground; a molding step (step S22) in which a molded body is prepared by molding the powder of the ground rare earth sintered magnet alloy; a sintering step (step S23) in which the molded body is sintered at a predetermined sintering temperature to obtain a sintered body; an aging step (step S24) in which the sintered body is aged to enhance the magnetic properties of the rare earth sintered magnet 1, such as coercivity; and a sintered body cooling step (step S25) in which the aged sintered body is cooled. Each step will be described below.

In the grinding step S21, a rare earth sintered magnet alloy having a (Nd,Pr,RH,La,Sm)-Fe-B crystalline phase and a (Nd,Pr,RH,La,Sm)-O phase, manufactured according to the rare earth sintered magnet alloy manufacturing process shown in Figure 6, is ground to obtain a rare earth sintered magnet alloy powder with a particle size of 200 μm or less, preferably 0.5 μm to 100 μm, and further, considering magnetization performance, approximately 1 μm to 10 μm. The grinding of the rare earth sintered magnet alloy is carried out using, in one example, an agate mortar and pestle, a stamp mill, a jaw crusher, or a jet mill. In particular, when reducing the particle size of the powder, it is preferable to grind the rare earth sintered magnet alloy in an atmosphere containing an inert gas. Grinding the rare earth sintered magnet alloy in an atmosphere containing an inert gas can suppress the incorporation of oxygen into the powder. However, if the atmosphere during crushing does not affect the magnetic properties of the magnet, the crushing of the rare earth sintered magnet alloy may be carried out in the atmosphere. When manufacturing the rare earth sintered magnet 1 of Embodiments 1 and 2, the La and Sm of the rare earth sintered magnet alloy used when manufacturing the rare earth sintered magnet 1 of Embodiment 3 can be replaced with a rare earth element R other than Nd, Pr and the heavy rare earth element RH. In other words, a rare earth sintered magnet alloy having a (Nd, Pr, RH, R)-Fe-B crystalline phase and a (Nd, Pr, RH, R)-O phase should be crushed.

In the molding process of step S22, the powder of the rare-earth sintered magnet alloy is compressed and molded in a mold with a magnetic field applied to it to prepare a molded body. Here, the applied magnetic field can be 2T in one example. Note that molding may be performed without applying a magnetic field.

In the sintering step S23, the compression-molded body is held at a sintering temperature of 950°C to 1300°C, preferably 1000°C to less than 1150°C, for a period of 0.1 hours to 10 hours, preferably 1.0 hour to 6.0 hours, to obtain a sintered body. Sintering is preferably carried out in an atmosphere containing an inert gas or in a vacuum to suppress oxidation. Sintering may also be carried out while applying a magnetic field.

In the case of Figure 7, the aging process in step S24 includes the first aging process in step S24-1, the second aging process in step S24-2, the third aging process in step S24-3, and the fourth aging process in step S24-4. The aging is preferably carried out in an atmosphere containing an inert gas or in a vacuum to suppress oxidation.

In the first aging step of step S24-1, the obtained sintered body is held at a first aging temperature, which is below the sintering temperature, for a period of 0.1 hours to 10 hours, preferably 0.5 hours to 5 hours. Specifically, the first aging temperature is a temperature in the range of 700°C to less than 950°C, which is lower than the sintering temperature.

In the second aging step of step S24-2, after the first aging step, the sintered body held in the first aging step is held at a second aging temperature, which is lower than the first aging temperature, for a period of 0.1 hours to 10 hours, preferably 1.0 hour to 7 hours. Specifically, the second aging temperature is a temperature in the range of 450°C to less than 700°C, which is lower than the first aging temperature.

In the third aging step of step S24-3, after the second aging step, the sintered body held in the second aging step is raised again to the first aging temperature, specifically a temperature in the range of 700°C or more and less than 950°C, and held at the first aging temperature for a time in the range of 0.1 hours or more and 10 hours or less, preferably 0.5 hours or more and 5 hours or less.

In the fourth aging step of step S24-4, after the third aging step, the sintered body held in the third aging step is again held at the second aging temperature, specifically at a temperature in the range of 450°C to less than 700°C, for a period of 0.1 hours to 10 hours, preferably 1.0 hour to 7 hours.

Finally, in the sintered body cooling step S25, the sintered body held in the fourth aging step is held at a cooling temperature in the range of 200°C to less than 450°C for a period of 0.1 hours to 5 hours. Afterward, it is cooled to room temperature to produce the diffusion precursor for the rare earth sintered magnet 1. Cooling is preferably carried out in an atmosphere containing an inert gas or in a vacuum to suppress oxidation.

In this manner, a diffusion precursor, which is a sintered body having the final shape of the rare-earth sintered magnet 1, is formed.

Returning to Figure 5, in the grain boundary diffusion step of step S30, heat treatment is performed under conditions in which the diffusion precursor formed in step S25 and the heavy rare earth element RH are present, causing the heavy rare earth element RH to diffuse into the diffusion precursor at the grain boundaries. In one example, the heat treatment is performed at a temperature below the sintering temperature in the sintering step of step S23 to hold the diffusion precursor. The grain boundary diffusion step may be performed simultaneously with the aging step of step S24. In the grain boundary diffusion step, the heavy rare earth element RH is selectively diffused into at least a portion of the outer edge of the Sm-enriched portion 41 of the first subphase 21 and uniformly diffused into the second subphase 22. Known grain boundary diffusion methods can be used for the treatment in the grain boundary diffusion step. Various techniques have been proposed for grain boundary diffusion methods depending on the supply form of the heavy rare earth element RH, with coating diffusion, sputtering diffusion, and vapor diffusion being typical methods. These typical grain boundary diffusion methods will be described below.

<Application and Diffusion Method>
In the coating diffusion method, the grain boundary diffusion step includes a diffusion element attachment step in which a heavy rare earth element supply unit, which is a material containing heavy rare earth elements RH and serves as a source of heavy rare earth elements RH to the diffusion precursor, is attached, and a diffusion heat treatment step in which heat treatment is performed to diffuse the heavy rare earth elements RH from the heavy rare earth element supply unit to the diffusion precursor. In the diffusion element attachment step, a slurry obtained by mixing powdered heavy rare earth element mixture with water or an organic solvent is attached to the surface of the diffusion precursor. The slurry attached to the surface of the diffusion precursor becomes the heavy rare earth element supply unit. The slurry can be attached by spray spraying, dip coating, spin coating, screen printing, electrodeposition, etc. In the diffusion heat treatment step, the diffusion precursor to which the heavy rare earth element supply unit is attached is heat-treated at a diffusion temperature below the sintering temperature in the sintering step of step S23, thereby diffusing the heavy rare earth elements RH into the interior of the diffusion precursor. The heat treatment conditions are a diffusion temperature below the sintering temperature and a time within the range of 0.1 hours to 100 hours. The diffusion temperature is, in one example, in the range of 300°C to 1000°C, and is lower than the sintering temperature. The heat treatment is preferably carried out in an atmosphere containing an inert gas or in a vacuum to suppress oxidation.

<Sputter Diffusion Method>
Similar to the coating diffusion method, the sputtering diffusion method includes a grain boundary diffusion step, which comprises a diffusion element deposition step and a diffusion heat treatment step. In the diffusion element deposition step, a thin film of a single metal or alloy composition of heavy rare earth element RH is formed on the surface of the diffusion precursor in a dry environment. The thin film formed on the surface of the diffusion precursor becomes a heavy rare earth element supply section. In one example, the thin film is formed by the sputtering method. In the diffusion heat treatment step, the diffusion precursor on which the heavy rare earth element supply section has been formed is heat-treated at a diffusion temperature below the sintering temperature in the sintering step of step S23, thereby diffusing the heavy rare earth element RH into the interior of the diffusion precursor. The heat treatment conditions are a diffusion temperature below the sintering temperature and a time within the range of 0.1 hours to 100 hours. In one example, the diffusion temperature is within the range of 300°C to 1000°C, which is lower than the sintering temperature. The heat treatment is preferably performed in an atmosphere containing an inert gas or in a vacuum to suppress oxidation.

<Vapor Diffusion Method>
In the vapor diffusion method, after placing the diffusion precursor and the heavy rare earth element supply unit in a vacuum furnace, the diffusion precursor is heat-treated in the vacuum furnace at a temperature lower than the sintering temperature in the sintering process of step S23, thereby diffusing the heavy rare earth element RH into the interior of the diffusion precursor. In the heat treatment, the heavy rare earth element supply source is converted to a gas phase by vacuum heating, and the heavy rare earth element RH is supplied to the diffusion precursor through the gas phase. The conditions for the heat treatment are a diffusion temperature lower than the sintering temperature and a time within the range of 0.1 hours to 100 hours. In one example, the diffusion temperature is within the range of 600°C to 900°C, which is lower than the sintering temperature. Furthermore, in the vapor diffusion method, unlike the coating diffusion method and sputter diffusion method, it is not necessary to attach the heavy rare earth element supply unit to the diffusion precursor, and the diffusion element attachment process can be omitted, thus shortening the time of the grain boundary diffusion process.

Returning to Figure 5, in the final step S40, the cooling process, the diffusion precursor in which the heavy rare earth element RH has been diffused in the grain boundary diffusion process is held at a temperature of less than 200°C for a time in the range of 0.1 hours to 5 hours. After that, by cooling to room temperature, the rare earth sintered magnet 1 shown in Embodiments 1 to 3 is formed. In Embodiment 1, a rare earth sintered magnet 1 is formed in which the heavy rare earth element RH is present on at least a portion of the surface of the main phase 10 containing the heavy rare earth element RH. In Embodiment 2, a rare earth sintered magnet 1 is formed in which the heavy rare earth element RH is diffused in the subphase 20 existing between the main phase 10 containing the heavy rare earth element RH. In Embodiment 3, a rare earth sintered magnet 1 is formed comprising a main phase 10 containing the heavy rare earth element RH, a first subphase 21 in which the heavy rare earth element RH is diffused so as to selectively surround the outer edge of the Sm-enriched portion 41, and a second subphase 22 in which the heavy rare earth element RH is uniformly diffused. Cooling is preferably carried out in an atmosphere containing an inert gas or in a vacuum to suppress oxidation.

As described above, by diffusing heavy rare earth element RH between grain boundaries into a diffusion precursor having the final shape of the rare earth sintered magnet 1, a rare earth sintered magnet 1 of the desired shape can be obtained.

As described above, in Embodiment 4, a rare earth sintered magnet alloy powder is obtained by crushing a rare earth sintered magnet alloy having a (Nd,Pr,RH,La,Sm)-Fe-B crystalline phase and a (Nd,Pr,RH,La,Sm)-O phase. The molded body is then sintered to form a sintered body, and the sintered body is then aged to produce a rare earth sintered magnet 1. This makes it possible to produce a rare earth sintered magnet 1 having the structure described in Embodiment 3. Furthermore, since the rare earth sintered magnet alloy is produced after adjusting the heavy rare earth elements to a desired concentration, and the rare earth sintered magnet 1 is produced using this rare earth sintered magnet alloy, the heavy rare earth element RH can easily enter the interior of the main phase 10.

Furthermore, in Embodiment 4, the temperature and time in the sintering process, aging process, and sintered body cooling process are controlled. In particular, in the first aging process, the obtained sintered body is held at a first aging temperature, which is below the sintering temperature, for a range of 0.1 hours to 10 hours, preferably 0.5 hours to 5 hours. In the second aging process, the sintered body is held at a second aging temperature, which is below the first aging temperature, for a range of 0.1 hours to 10 hours, preferably 1.0 hour to 7 hours. In the third aging process, the temperature is raised again to the first aging temperature, and the sintered body is held at the first aging temperature for a range of 0.1 hours to 10 hours, preferably 0.5 hours to 5 hours. In the fourth aging process, the sintered body is again held at the second aging temperature for a range of 0.1 hours to 10 hours, preferably 1.0 hour to 7 hours. In this way, the temperature and time are controlled so that the first aging process and the second aging process are performed in two sets. This creates a state in which the sintered body is repeatedly held in a temperature range with unstable energy states. As a result, a first main phase 11 consisting of CNd > CPr and a second main phase 12 consisting of CNd < CPr coexist, making it possible to make the concentration of heavy rare earth element RH in the first main phase 11 higher than the concentration of heavy rare earth element RH in the second main phase 12. In other words, the rare earth sintered magnet 1 has two types of main phases 10, the first main phase 11 and the second main phase 12. Focusing on the core portions 11c and 12c of the two types of main phases 10, it is possible to selectively manufacture a rare earth sintered magnet 1 in which the sum of the concentration of Nd and the concentration of heavy rare earth element RH in the first main phase 11 is higher than the concentration of Pr, and conversely, the concentration of Pr in the second main phase 12 is higher than the sum of the concentration of Nd and the concentration of heavy rare earth element RH.

Furthermore, in addition to the first main phase 11 and the second main phase 12 in Embodiment 1, a rare earth sintered magnet 1 having a characteristic microstructure is selectively manufactured, comprising a crystalline first subphase 21 based on an oxide phase whose main component is (Nd, Pr, RH, La, Sm)-O, and a crystalline second subphase 22 whose main component is (Nd, Pr, RH, La)-O, wherein the concentration of Sm is higher in the first subphase 21 than in the second subphase 22, and an Sm-enriched portion 41 is formed in the first subphase 21.

Furthermore, in the manufacturing method of the rare earth sintered magnet 1 according to Embodiment 4, an R-Fe-B type rare earth sintered magnet alloy containing Nd and Pr as rare earth element R is pulverized, a molded body of the R-Fe-B type rare earth sintered magnet alloy powder is sintered, and an aging treatment is performed to form a diffusion precursor having a first main phase 11 and a second main phase 12, wherein the concentration of heavy rare earth element RH in the first main phase 11 is higher than the concentration of heavy rare earth element RH in the second main phase 12. In the manufacturing method of the rare earth sintered magnet 1 according to Embodiment 4, by heat treatment that diffuses heavy rare earth element RH between grain boundaries in the diffusion precursor, it is possible to manufacture a rare earth sintered magnet 1 in which heavy rare earth element RH is present inside the first main phase 11 and the second main phase 12, as well as on a part of the surface of the first main phase 11 and the second main phase 12, or a rare earth sintered magnet 1 in which heavy rare earth element RH is present in the sub-phase 20.

Furthermore, in the method for manufacturing the rare earth sintered magnet 1 according to Embodiment 4, a R-Fe-B type rare earth sintered magnet alloy containing Nd, Pr, La, and Sm as rare earth element R is pulverized, a molded body of the R-Fe-B type rare earth sintered magnet alloy powder is sintered, and an aging treatment is performed to form a diffusion precursor having a first main phase 11 containing the heavy rare earth element RH, a second main phase 12 having a lower concentration of the heavy rare earth element RH than the first main phase 11, a first sub-phase 21 having an Sm-enriched portion 41, and a second sub-phase 22 having a lower Sm concentration than the first sub-phase 21. In the manufacturing method of the rare earth sintered magnet 1 according to Embodiment 4, by heat treatment to diffuse heavy rare earth elements RH across grain boundaries in the diffusion precursor, a rare earth sintered magnet 1 can be produced in which, in the first sub-phase 21, the outer periphery of the Sm-enriched portion 41 is selectively surrounded by heavy rare earth elements RH, and in the second sub-phase 22, the heavy rare earth elements RH are uniformly distributed.

Compared to the R-T-B sintered magnet described in Patent Document 1, the R-T-B sintered magnet described in Patent Document 1 has one or more rare earth elements, with heavy rare earth element RH being essential, and has one type of main phase particle composed of a core portion and a shell portion. In other words, in Patent Document 1, all main phase particles contain the heavy rare earth element RH. On the other hand, the main phase 10 of the rare earth sintered magnet 1 according to Embodiment 1 is a mixture of a first main phase 11 in which the heavy rare earth element RH is contained in the core portion 11c, and a second main phase 12 in which the heavy rare earth element RH is hardly contained in the core portion 12c. In other words, the heavy rare earth element RH is selectively arranged in the first main phase 11 of the two types of main phases 10. Thus, compared to the R-T-B sintered magnet described in Patent Document 1, which requires that the heavy rare earth element RH be contained in all of the single type of main phase, the rare earth sintered magnet 1 according to Embodiment 1, which only requires that the heavy rare earth element RH be contained in one of the two types of main phases 10 (the first main phase 11), can reduce the amount of heavy rare earth element RH used. Furthermore, since the volume of the subphase 20 in relation to the total volume of the rare earth sintered magnet 1 is extremely small, even if the heavy rare earth element RH is diffused and present in the subphase 20, the amount of heavy rare earth element RH used can be reduced compared to the technology in Patent Document 1.

Furthermore, compared to the technology described in Patent Document 2, in order to obtain magnetic properties equivalent to those of the rare earth sintered magnet 1 according to Embodiment 1 using the method described in Patent Document 2, a large amount of heavy rare earth element RH must be added, as will be described later. In other words, when trying to obtain the same magnetic properties, the rare earth sintered magnet 1 according to Embodiment 1 can use less heavy rare earth element RH compared to the technology described in Patent Document 2. Also, the method described in Patent Document 2 includes hot working, but the manufacturing method of the rare earth sintered magnet 1 according to Embodiment 4 does not include hot working. Therefore, the reduction in particle size of the main phase 10 is suppressed, and the decrease in residual magnetic flux density and magnetization can be suppressed compared to the rare earth magnet manufactured using the technology described in Patent Document 2.

This makes it possible to obtain a rare-earth sintered magnet 1 that improves magnetic properties compared to conventional methods while suppressing the use of heavy rare-earth elements (RH) compared to conventional methods.

Embodiment 5.
Embodiment 5 describes a rotor using the rare earth sintered magnet 1 from Embodiments 1, 2, and 3, manufactured by the manufacturing method of Embodiment 4. Figure 8 is a schematic cross-sectional view showing an example of the configuration of a rotor equipped with a rare earth sintered magnet according to Embodiment 5. Figure 8 shows a cross-section of the rotor 100 in a direction perpendicular to the rotation axis RA.

The rotor 100 is rotatable around the rotation axis RA. The rotor 100 comprises a rotor core 101 and rare earth sintered magnets 1 inserted into magnet insertion holes 102 provided in the rotor core 101 along the circumferential direction of the rotor 100. Figure 8 shows an example in which four magnet insertion holes 102 are provided in the rotor core 101 and four rare earth sintered magnets 1 are inserted into the magnet insertion holes 102, but the number of magnet insertion holes 102 and rare earth sintered magnets 1 may be changed according to the design of the rotor 100. The rotor core 101 is formed by stacking multiple disc-shaped electromagnetic steel sheets in the axial direction of the rotation axis RA.

The rare earth sintered magnets 1 were manufactured according to the manufacturing method described in Embodiment 4. The four rare earth sintered magnets 1 are each inserted into their corresponding magnet insertion holes 102. The four rare earth sintered magnets 1 are magnetized such that the magnetic poles of the rare earth sintered magnets 1 on the radially outer side of the rotor 100 differ from those of adjacent rare earth sintered magnets 1.

Thus, the rotor 100 according to Embodiment 5 is equipped with a rare-earth sintered magnet 1 according to Embodiment 1, Embodiment 2, or Embodiment 3, which can achieve improved magnetic properties at room temperature and suppression of the decrease in magnetic properties with increasing temperature. In this way, the rare-earth sintered magnet 1 suppresses the decrease in magnetic properties with increasing temperature while suppressing the use of heavy rare-earth elements RH compared to conventional methods and maintaining high residual magnetic flux density and coercivity. Therefore, the decrease in magnetic properties is suppressed even in high-temperature environments exceeding 100°C. This makes it possible to improve magnetic properties and magnetization while replacing expensive, geographically unevenly distributed, and procurement-risky Nd and heavy rare-earth elements RH with inexpensive rare-earth elements, and to stabilize the operation of the rotor 100 even in high-temperature environments exceeding 100°C. Furthermore, since the rare-earth sintered magnet 1 according to Embodiment 1, Embodiment 2, or Embodiment 3 has superior magnetization performance compared to conventional methods, it is possible to magnetize the rotor 100 in an assembled state with the rare-earth sintered magnet 1 set, thus simplifying the manufacturing process. Furthermore, since a magnetization process with reduced voltage can be achieved, it also contributes to energy saving.

Embodiment 6.
Embodiment 6 describes a rotating machine equipped with the rotor 100 from Embodiment 5. Figure 9 is a schematic cross-sectional view showing an example of the configuration of the rotating machine according to Embodiment 6. Figure 9 shows a cross-section perpendicular to the rotation axis RA of the rotor 100.

The rotating machine 120 comprises a rotor 100, as described in Embodiment 5, which is rotatable around a rotation axis RA, and an annular stator 130 provided coaxially with the rotor 100 and positioned opposite the rotor 100. The stator 130 is formed by stacking multiple electromagnetic steel sheets in the axial direction of the rotation axis RA. The configuration of the stator 130 is not limited to this, and existing configurations can also be adopted. The stator 130 has teeth 131 that protrude toward the rotor 100, provided along the inner surface of the stator 130. Windings 132 are attached to the teeth 131. In one example, the winding method of the windings 132 may be concentrated winding or distributed winding. In other words, the stator 130 has windings 132 attached to teeth 131 that protrude toward the rotor 100 on the inner surface on the side where the rotor 100 is located, and has an annular structure positioned opposite the rotor 100. The rotor 100 inside the rotating machine 120 has two or more magnetic poles; that is, it is sufficient to have two or more rare-earth sintered magnets 1. Also, although Figure 9 shows an example of a magnet-embedded rotor 100, a surface-magnet type rotor 100 in which rare-earth sintered magnets 1 are fixed to the outer circumference with adhesive is also acceptable.

Thus, the rotating machine 120 in Embodiment 6 is equipped with a rare earth sintered magnet 1 according to Embodiment 1, Embodiment 2, or Embodiment 3, which can achieve improved magnetic properties at room temperature and suppression of the decrease in magnetic properties with increasing temperature. In this way, the rare earth sintered magnet 1 suppresses the decrease in magnetic properties with increasing temperature while suppressing the use of heavy rare earth elements RH compared to conventional methods and maintaining high residual magnetic flux density and coercivity. As a result, the decrease in magnetic properties is suppressed even in high-temperature environments exceeding 100°C. As a result, while substituting Nd and heavy rare earth elements RH, which are expensive, have high regional distribution and procurement risks, with inexpensive rare earth elements, magnetic properties and magnetization are improved, and the rotor 100 can be stably driven and the operation of the rotating machine 120 can be stabilized even in high-temperature environments exceeding 100°C.

The details of the rare-earth sintered magnet 1 of this disclosure will be described below with reference to examples and comparative examples.

In Examples 1 to 8, rare earth sintered magnets 1 are manufactured using the method shown in Embodiment 4, with samples of multiple rare earth sintered magnet alloys with different compositions represented as (Nd, Pr, Dy, La, Sm)-Fe-B. In Examples 1 to 8, a diffusion precursor is formed using rare earth sintered magnet alloys with varying Nd, Pr, Dy, La, and Sm content, and the rare earth sintered magnets 1 are manufactured by grain boundary diffusion of Tb, a heavy rare earth element RH, into the diffusion precursor so that Tb is 0.10 at.%. In other words, in Examples 1 to 8, rare earth sintered magnets 1 are manufactured from rare earth sintered magnet alloys represented as (Nd, Pr, Dy, La, Sm)-Fe-B, using the manufacturing method shown in Embodiment 4, with 0.10 at.% of the heavy rare earth element RH Tb diffused into them.

In Comparative Examples 1 to 14, rare earth sintered magnets 1 were experimentally produced by grain boundary diffusion of the heavy rare earth element RH, Tb, using samples represented by multiple rare earth sintered magnet alloys R-Fe-B with different compositions, using a general rare earth magnet manufacturing method as described in Patent Document 1 or Patent Document 2. In the samples of rare earth sintered magnets 1 from Comparative Examples 1 to 14, the R portion was changed.

In Comparative Examples 1 to 7, a rare earth sintered magnet 1 is produced using the manufacturing method described in Patent Document 1, in which 0.15 at.% of the heavy rare earth element RH, Tb, is diffused, either from a rare earth sintered magnet alloy where R is Nd, or from a rare earth sintered magnet alloy containing R = Nd and one or more elements selected from the group consisting of Dy, Pr, La, and Sm.

In Comparative Examples 8 to 14, a rare earth sintered magnet 1 is produced using the manufacturing method described in Patent Document 2, in which 0.15 at.% of the heavy rare earth element RH, Tb, is diffused, either from a rare earth sintered magnet alloy where R is Nd, or from a rare earth sintered magnet alloy containing R = Nd and one or more elements selected from the group consisting of Dy, Pr, La, and Sm.

Table 3 shows the general formulas, elemental content, microstructure analysis results, and magnetic properties and magnetization performance of rare earth sintered magnets for the examples and comparative examples. Table 3 shows the general formulas of the main phase 10 for each sample, which is the rare earth sintered magnet 1 of Examples 1 to 8 and Comparative Examples 1 to 14.

Next, a method for analyzing the microstructure of the rare earth sintered magnets 1 of Examples 1 to 8 and Comparative Examples 1 to 14 will be described. The microstructure of the rare earth sintered magnets 1 is determined by elemental analysis using a scanning electron microscope (SEM) and EPMA. Here, an FE-EPMA (manufactured by JEOL Ltd., product name: JXA-8530F) is used as the SEM and EPMA. The conditions for elemental analysis are an acceleration voltage of 15.0 kV, an irradiation current of 2.271 ^e⁻⁰⁸ A, an irradiation time of 130 ms, a pixel count of 512 pixels × 512 pixels, a magnification of 5000x, and one integration cycle.

Next, the method for evaluating the magnetic properties of the rare-earth sintered magnets 1 of Examples 1 to 8 and Comparative Examples 1 to 14 will be described. The magnetic properties are evaluated by measuring the coercivity of multiple samples using a pulse-excitation type BH tracer. The maximum applied magnetic field by the BH tracer is 6T or higher, which is the state in which the rare-earth sintered magnet 1 is fully magnetized. In addition to the pulse-excitation type BH tracer, any device that can generate a maximum applied magnetic field of 6T or higher may be used, such as a DC recording magnetometer (also called a DC BH tracer), a vibrating sample magnetometer (VSM), a magnetic property measurement system (MPMS), or a physical property measurement system (PPMS). Measurements are performed in an atmosphere containing an inert gas such as nitrogen. The magnetic properties of each sample are measured by detecting the magnetization picked up by a search coil or magnetic sensor from the rare-earth sintered magnet 1, which has been magnetized by the applied magnetic field. The magnetic properties are measured from the J-H curve or B-H curve, which are the measured magnetic hysteresis curves. The magnetic properties of each sample are measured at two different temperatures, a first measurement temperature T1 and a second measurement temperature T2. The temperature coefficient α [%/°C] of the remanent magnetic flux density is the ratio of the difference between the remanent magnetic flux density at the first measurement temperature T1 and the remanent magnetic flux density at the second measurement temperature T2, divided by the temperature difference (T2-T1). The temperature coefficient β [%/°C] of the coercivity is the ratio of the difference between the coercivity at the first measurement temperature T1 and the coercivity at the second measurement temperature T2, divided by the temperature difference (T2-T1). Therefore, the smaller the absolute values of the temperature coefficients of the magnetic properties, |α| and |β|, the more the decrease in the magnetic properties of the magnet with respect to temperature rise is suppressed.

Furthermore, the magnetization performance is measured by calculating the susceptibility to magnetization from the ratio of the magnetic flux density measured from the magnetic hysteresis observed when an arbitrary magnetic field is applied at a constant permeance coefficient to the magnetic flux density measured from the magnetic hysteresis observed when a saturating magnetic field is applied. If a high susceptibility to magnetization can be obtained even at lower magnetic fields, it can be said that the magnetization performance is high.

First, the analytical results for each sample from Examples 1 to 8 and Comparative Examples 1 to 14 will be explained. Although not shown in the diagram, elemental mapping of each sample from Examples 1 to 8 using FE-EPMA revealed that, with RH set to Dy and Tb, and R set to one or more rare earth elements selected from elements other than Nd, Pr, and RH, the main phase 10 containing crystal grains that satisfy the general formula (Nd, Pr, Dy, Tb, R)-Fe-B and are based on the Nd ₂ Fe ₁₄ B crystal structure, has core portions 11c, 12c and shell portions 11s, 12s covering the core portions 11c, 12c. Furthermore, it was confirmed that the main phase 10 contains a mixture of a first main phase 11 where CNd > CPr and a second main phase 12 where CNd < CPr. Furthermore, it can be confirmed that when the concentration of Dy in the first main phase 11 is C1Dy and the concentration of Dy in the second main phase 12 is C2Dy, C1Dy > C2Dy.

Here, the concentration difference shown for "the first main phase 11 where CNd > CPr and the second main phase 12 where CNd < CPr" means that a clear difference is observed between the detection intensity of Nd and the detection intensity of Pr, as determined by mapping analysis using EPMA. Specifically, taking the first main phase 11 as an example, the concentration of Nd in the core portion 11c is detected by EPMA at a higher-than-average level, while the concentration of Pr is detected by EPMA at a level near the lower limit. The second main phase 12 is the opposite of the first main phase 11.

Furthermore, the concentration difference indicated by "C1Dy > C2Dy" means that, as determined by mapping analysis using EPMA, there is a clear difference between the detection intensity of Dy in the first main phase 11 and the detection intensity of Dy in the second main phase 12. Specifically, the detection intensity of Dy in the first main phase 11 is higher than average according to EPMA, while the detection intensity of Dy in the second main phase 12 is near the lower limit according to EPMA.

Furthermore, when R = La, Sm, it can be confirmed that the rare earth sintered magnet 1, in addition to the first main phase 11 and second main phase 12 in Embodiment 1, also has a crystalline first sub-phase 21 based on an oxide phase whose main component is represented as (Nd, Pr, Dy, Tb, La, Sm)-O, and a crystalline second sub-phase 22 whose main component is represented as (Nd, Pr, Dy, Tb, La)-O. Moreover, it can be confirmed that the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22.

Table 3 shows that, if the sum of the concentrations of Nd and the heavy rare earth element Tb is denoted as C(Nd, Tb), then for samples where the state of the first main phase 11 (C(Nd, RH) > CPr) and the second main phase 12 (C(Nd, RH) < CPr) was confirmed, "○" is entered in the columns for the first main phase 11 and the second main phase 12, respectively. For samples where this could not be confirmed, "×" is entered in the columns for the first main phase 11 and the second main phase 12, respectively. The inequality sign indicates a concentration difference that clearly shows a difference between the detection intensity of Nd and Dy and the detection intensity of Pr. To give a specific example, in the case of the first main phase 11, the concentrations of Nd and Dy indicate that the detection intensity of EPMA is higher than average, while the concentration of Pr indicates that the detection intensity of EPMA is near the lower limit. In the case of the second main phase 12, the opposite is true compared to the case of the first main phase 11. If only C(Nd,RH) < CPr, as in the second main phase 12, is confirmed, then "○" is entered only in the second main phase 12 column, and "×" is entered in the first main phase 11 column.

Furthermore, Table 3 shows that for samples having a crystalline first subphase 21 based on an oxide phase whose main component is represented as (Nd, Pr, Dy, Tb, La, Sm)-O, and a crystalline second subphase 22 whose main component is represented as (Nd, Pr, Dy, Tb, La)-O, and where it was confirmed that the concentration of Sm is higher in the first subphase 21 than in the second subphase 22, "○" is entered in both the first subphase 21 and second subphase 22 columns. For samples where this could not be confirmed, "×" is entered in both the first subphase 21 and second subphase 22 columns. Also, for samples where only one subphase 20 exists, or where there is no difference in Sm concentration between the subphases 20, it is assumed that only the first subphase 21 exists, and "○" is entered only in the first subphase 21 column, while "×" is entered in the second subphase 22 column. Furthermore, the concentration difference between the first subphase 21 and the second subphase 22 indicates, based on mapping analysis using EPMA, that the detection intensity of Sm is, on average, higher in the first subphase 21 than in the second subphase 22.

Furthermore, from the intensity ratio of the elemental mapping obtained by analysis with FE-EPMA, it can be confirmed that there are more instances of the first main phase 11 where C(Nd,RH) > CPr than instances of the second main phase 12 where C(Nd,RH) < CPr. When focusing on the shell parts 11s and 12s of the core-shell structure, with the concentration of Dy in the core parts 11c and 12c being CDy and the concentration of Dy in the shell parts 11s and 12s being SDy, it can be confirmed that the first main phase 11 satisfies the relationships CNd > SNd, CPr < SPr, CDy > SDy, and the second main phase 12 satisfies the relationships CNd < SNd, CPr > SPr, CDy < SDy.

Next, the measurement results of the magnetic properties of each sample in Examples 1 to 8 and Comparative Examples 1 to 14 will be described. The shape of each sample used for magnetic measurement is a block shape with length, width, and height all of 7 mm. The first measurement temperature T1 is 23°C, and the second measurement temperature T2 is 200°C. 23°C is room temperature. The second measurement temperature T2 of 200°C is a temperature that can occur in the operating environment of automotive motors and industrial motors.

First, the residual magnetic flux density and coercivity of each sample in Examples 1 to 8 and Comparative Examples 2 to 14 are determined in comparison with Comparative Example 1. If the residual magnetic flux density and coercivity values of each sample at 23°C are within 1% of the values in Comparative Example 1, which is considered to be within the measurement error, it is judged as "equivalent." If the values are 1% or more higher, it is judged as "good." If the values are 1% or less lower, it is judged as "poor."

Next, the temperature coefficient α of the remanent magnetic flux density is calculated using the remanent magnetic flux density at 23°C (first measurement temperature T1) and the remanent magnetic flux density at 200°C (second measurement temperature T2). Similarly, the temperature coefficient β of the coercivity is calculated using the coercivity at 23°C (first measurement temperature T1) and the coercivity at 200°C (second measurement temperature T2). The temperature coefficients of remanent magnetic flux density and coercivity for each sample from Examples 1 to 8 and Comparative Examples 2 to 14 are judged by comparing them with Comparative Example 1. For each sample, if the values are within ±1% of the absolute value of the temperature coefficient of remanent magnetic flux density |α| and the absolute value of the temperature coefficient of coercivity |β| compared with the sample from Comparative Example 1, it is judged as "equivalent" if the values are within ±1% of what is considered a measurement error; if the values are lower than -1%, it is judged as "good"; and if the values are higher than +1%, it is judged as "poor". For samples judged as "good," the temperature coefficient is smaller, which suppresses the decrease in magnetic properties with increasing temperature. This makes it possible to provide a rare-earth sintered magnet 1 that has stable magnetic properties even in high-temperature environments.

Next, the magnetization performance is calculated from the ratio of the magnetic flux density at the intersection of the magnetic hysteresis and permeance coefficient Pc at an applied magnetic field of 20 kOe to the magnetic flux density at the intersection of the magnetic hysteresis and permeance coefficient Pc at an applied magnetic field of 80 kOe, which is the saturated magnetization state. The magnetization performance of each sample from Examples 1 to 8 and Comparative Examples 2 to 14 is judged in comparison to Comparative Example 1. In other words, for each sample, if the magnetization performance is 1% or more compared to the sample from Comparative Example 1, which is considered to be a measurement error, it is judged as "equivalent or better," and if the value is lower than 1%, it is judged as "poor." For samples judged as "equivalent or better," a rare-earth sintered magnet 1 with high magnetization performance can be provided.

The results for determining the remanent magnetic flux density, coercivity, temperature coefficient of remanent magnetic flux density, temperature coefficient of coercivity, and magnetization performance are shown in Table 3.

Comparative Example 1 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to the manufacturing method described in Patent Document 1, using Nd, Dy, Fe, and FeB as raw materials so that the composition is (Nd, Dy)-Fe-B. When the microstructure of this sample is observed according to the method described above, since Pr, La, and Sm are not added, the core-shell structure in the main phase 10 cannot be confirmed, and it cannot be confirmed that the concentration of Sm in the sub-phase 20 is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the remanent magnetic flux density is 1.25 T and the coercivity is 1900 kA/m. The temperature coefficients of the remanent magnetic flux density and coercivity are |α| = 0.185%/°C and |β| = 0.455%/°C, respectively. The magnetic attachment rate is 98.6%. These values for Comparative Example 1 are used as a reference.

Comparative Example 2 is a sample of a rare earth sintered magnet 1 with 0.15 at.% Tb diffused into it, prepared according to the manufacturing method described in Patent Document 1, using Nd, Fe, and FeB as raw materials so that it is Nd-Fe-B. When the microstructure of this sample is observed according to the method described above, since Pr, La, and Sm are not added, the core-shell structure in the main phase 10 cannot be confirmed, and it cannot be confirmed that the concentration of Sm in the sub-phase 20 is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, since the heavy rare earth elements RH, Dy and Tb, are not present in the main phase 10, the residual magnetic flux density is "good" and the coercivity is "poor". Also, the temperature coefficient of the residual magnetic flux density is "equivalent", the temperature coefficient of the coercivity is "equivalent", and the magnetization performance is "equivalent or better".

Comparative Example 3 is a sample of a rare earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to the manufacturing method described in Patent Document 1, using Nd, Pr, Fe, and FeB as raw materials so that it is (Nd, Pr)-Fe-B. When the microstructure of this sample is observed according to the method described above, although the main phase 10, which is a mixture of Nd and Pr, can be confirmed due to the addition of Pr, a core-shell structure is not formed. Also, since La and Sm are not added, it cannot be confirmed that the concentration of Sm in the sub-phase 20 is higher in the first sub-phase 21 than in the second sub-phase 22. When the magnetic properties of this sample are evaluated according to the method described above, the residual magnetic flux density is "good" and the coercivity is "poor" because the heavy rare earth elements RH, Dy and Tb, are not present in the base material. Also, the temperature coefficient of coercivity is "poor" due to the addition of Pr. Since the manufacturing method does not involve hot working, the magnetization performance will be "equivalent or better." Furthermore, the temperature coefficient of the residual magnetic flux density will be "equivalent." In addition, diffusing Tb, which is a heavy rare earth element RH, into a diffusion precursor, which is a base material that does not contain heavy rare earth elements RH, will not improve the magnetic properties.

Comparative Example 4 is a sample of a rare earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to the manufacturing method described in Patent Document 1, using Nd, Pr, Dy, Fe, and FeB as raw materials so that the composition is (Nd, Pr, Dy)-Fe-B. When the microstructure of this sample is observed according to the method described above, the main phase 10, which is a mixture of Nd and Pr, can be confirmed, but it does not form a core-shell structure. Also, since La and Sm are not added, it cannot be confirmed that the concentration of Sm in the subphase 20 is higher in the first subphase 21 than in the second subphase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the residual magnetic flux density is "equivalent" because the addition of Pr and the addition of Dy, a heavy rare earth element RH, is the same as in Comparative Example 1. The addition of Pr, in addition to the heavy rare earth element Dy, results in "good" coercivity and a "poor" temperature coefficient of coercivity. Since the manufacturing method does not include hot working, the magnetization performance is "equal to or better." The temperature coefficient of residual magnetic flux density is "equal," which reflects the fact that although the addition of Tb and Pr increases the magnetic anisotropy of the main phase 10 and improves coercivity, it does not result in the optimal microstructure of the main phase 10 and subphase 20.

Comparative Example 5 is a sample of a rare earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to the manufacturing method described in Patent Document 1, using Nd, La, Sm, Fe, and FeB as raw materials so that the composition is (Nd, La, Sm)-Fe-B. When the microstructure of this sample is observed according to the method described above, the core-shell structure of the main phase 10 cannot be confirmed because Pr is not added. Also, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, the second sub-phase 22 does not exist. Furthermore, it cannot be confirmed that the concentration of Sm in the first sub-phase 21 is higher than that of the second sub-phase 22. In addition, when the magnetic properties of this sample are evaluated according to the method described above, the residual magnetic flux density is "good" and the coercivity is "poor" because the heavy rare earth element RH, Dy, is not present in the base material. Furthermore, the addition of La and Sm results in a "good" temperature coefficient of coercivity. Since the manufacturing method does not involve hot working, the magnetization performance is "equal to or better." The temperature coefficient of residual magnetic flux density is "equal." This is because, although the presence of La and Sm in the main phase 10 or sub-phase 20 results in a good temperature coefficient of magnetic properties, the magnetic properties at room temperature do not improve, reflecting that the microstructure of the main phase 10 and sub-phase 20 is not optimal. In addition, since the magnetic properties depend on the microstructure of the diffusion precursor, diffusing the heavy rare earth element RH, Tb, into such a diffusion precursor does not improve the magnetic properties.

Comparative Example 6 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused into it, prepared according to the manufacturing method described in Patent Document 1, using Nd, La, Sm, Fe, and FeB as raw materials so that the composition is (Nd, La, Sm)-Fe-B. The composition ratio of Nd, La, and Sm is different from that of Comparative Example 5. When the microstructure of this sample is observed according to the method described above, the core-shell structure of the main phase 10 cannot be confirmed because Pr is not added. Furthermore, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, the second sub-phase 22 does not exist. Moreover, it cannot be confirmed that the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the residual magnetic flux density is "good" and the coercivity is "poor" because the heavy rare earth element RH, Dy, is not present in the base material. However, by optimizing the amounts of La and Sm added, the temperature coefficient of the residual magnetic flux density becomes "good" and the temperature coefficient of the coercivity becomes "good". Since the manufacturing method does not include hot working, the magnetization performance is "equal to or better". This is because, although the temperature coefficient of the magnetic properties shows good results due to the presence of La and Sm in the main phase 10 or sub-phase 20, the magnetic properties at room temperature do not improve because the heavy rare earth element RH, Dy, is not added to the base material, reflecting that the microstructure of the main phase 10 and sub-phase 20 is not optimal. Even when the composition ratio of Nd, La, and Sm is changed, results almost the same as in Comparative Example 5 are obtained. Furthermore, since the magnetic properties depend on the microstructure of the diffusion precursor, diffusing the heavy rare earth element RH, Tb, into such a diffusion precursor does not improve its magnetic properties.

Comparative Example 7 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to the manufacturing method described in Patent Document 1, using Nd, Pr, La, Sm, Fe, and FeB as raw materials such as (Nd, Pr, La, Sm)-Fe-B. When the microstructure of this sample is observed according to the method described above, although a main phase 10 in which Nd and Pr are mixed can be confirmed due to the addition of Pr, it does not form a core-shell structure. Furthermore, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, a second sub-phase 22 does not exist. Moreover, it cannot be confirmed that the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the residual magnetic flux density is "good" and the coercivity is "poor" because the heavy rare earth element RH, Dy, is not present in the base material. In addition, while the temperature coefficients of the residual magnetic flux density and coercivity should be "good" by optimizing the amounts of La and Sm added, the addition of Pr has reduced only the temperature coefficient of coercivity to "equivalent". Since the manufacturing method does not include hot working, the magnetization performance is "equivalent or better".

Comparative Example 8 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, manufactured according to a manufacturing method including hot working described in Patent Document 2, using Nd, Dy, Fe, and FeB as raw materials so that the composition is (Nd, Dy)-Fe-B. When the microstructure of this sample is observed according to the method described above, since Pr, La, and Sm are not added, the core-shell structure in the main phase 10 cannot be confirmed, and the concentration of Sm in the sub-phase 20 cannot be confirmed as being higher in the first sub-phase 21 than in the second sub-phase 22. However, the refinement of the microstructure, which is a characteristic of magnets manufactured by hot working, can be confirmed. When the magnetic properties of this sample are evaluated according to the method described above, the coercivity is "good" due to the refinement of the magnetic powder, and the temperature coefficient of coercivity is "equivalent". Also, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of residual magnetic flux density is "equivalent". This result reflects the fact that while the absolute value of coercivity and the temperature coefficient of coercivity improve with the refinement of magnetic powder through hot working, the magnetic moments are less likely to align, leading to a decrease in residual magnetic flux density and deterioration of magnetization performance. Furthermore, since magnetic properties depend on the microstructure of the diffusion precursor, diffusing heavy rare earth element RH, such as Tb, into such a diffusion precursor does not improve its magnetic properties.

Comparative Example 9 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, manufactured using Nd, Fe, and FeB as raw materials to form Nd-Fe-B, according to a manufacturing method including hot working described in Patent Document 2. When the microstructure of this sample is observed according to the method described above, since Pr, La, and Sm are not added, the core-shell structure in the main phase 10 cannot be confirmed, and the concentration of Sm in the sub-phase 20 cannot be confirmed as being higher in the first sub-phase 21 than in the second sub-phase 22. However, the refinement of the microstructure, which is a characteristic of magnets manufactured by hot working, can be confirmed. When the magnetic properties of this sample are evaluated according to the method described above, the coercivity is "good" due to the refinement of the magnetic powder, and the temperature coefficient of coercivity is "equivalent". Also, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of residual magnetic flux density is "equivalent". This result reflects a decrease in residual magnetic flux density, although coercivity has improved due to the refinement of the microstructure through hot working. Furthermore, since magnetic properties depend on the microstructure of the diffusion precursor, diffusing the heavy rare earth element RH, Tb, into such a diffusion precursor does not improve its magnetic properties.

Comparative Example 10 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, manufactured according to a manufacturing method including hot working described in Patent Document 2, using Nd, Pr, Fe, and FeB as raw materials so that the composition is (Nd, Pr)-Fe-B. When the microstructure of this sample is observed according to the method described above, a core-shell structure is confirmed due to the addition of Pr and the hot working, but the core-shell structure is confirmed only in one type of main phase 10 where the Pr concentration in the core is high. Also, since La and Sm are not added, it cannot be confirmed that the concentration of Sm in the sub-phase 20 is higher in the first sub-phase 21 than in the second sub-phase 22. When the magnetic properties of this sample are evaluated according to the method described above, the coercivity is "good" and the temperature coefficient of coercivity is "equal" due to the refinement of the magnetic powder. Also, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of the residual magnetic flux density is "equivalent." This is because, although the coercivity is significantly improved to the level of rare-earth sintered magnet 1 with added Dy due to the formation of a core-shell structure with a high Pr concentration in the core, other properties reflect the refinement of the microstructure. Furthermore, since the magnetic properties depend on the microstructure of the diffusion precursor, which is the base material, diffusing Tb, a heavy rare-earth element RH, into such a diffusion precursor does not improve the magnetic properties.

Comparative Example 11 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, manufactured according to a manufacturing method including hot working described in Patent Document 2, using Nd, Pr, Dy, Fe, and FeB as raw materials so that the composition is (Nd, Pr, Dy)-Fe-B. When the microstructure of this sample is observed according to the method described above, a core-shell structure is confirmed due to the addition of Pr and the hot working, but the core-shell structure is confirmed only in one type of main phase 10 where the Pr concentration in the core is high. Also, since La and Sm are not added, it cannot be confirmed that the concentration of Sm in the sub-phase 20 is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the coercivity is "good" due to the refinement of the magnetic powder, and the temperature coefficient of coercivity is "equivalent". Also, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of the residual magnetic flux density is "equivalent." This is because, in addition to being produced by hot working, the coercivity is significantly improved due to the substitution of some of the highly anisotropic Dy with Nd, while other properties reflect the refinement of the microstructure due to hot working. Furthermore, since the magnetic properties depend on the microstructure of the diffusion precursor, diffusing the heavy rare earth element RH Tb into such a diffusion precursor does not improve the magnetic properties.

Comparative Example 12 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to a manufacturing method including hot working described in Patent Document 2, using Nd, La, Sm, Fe, and FeB as raw materials so that the composition is (Nd, La, Sm)-Fe-B. When the microstructure of this sample is observed according to the method described above, the core-shell structure of the main phase 10 cannot be confirmed because Pr is not added. Also, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, the second sub-phase 22 does not exist. Furthermore, it cannot be confirmed that the concentration of Sm in the first sub-phase 21 is higher than that of the second sub-phase 22. Also, when the magnetic properties of this sample are evaluated according to the method described above, the coercivity and temperature coefficient of coercivity are "good" due to the refinement of the magnetic powder. Furthermore, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of the residual magnetic flux density is "good". This is because, although the presence of La and Sm in the main phase 10 or sub-phase 20 results in a good temperature coefficient for the magnetic properties, the residual magnetic flux density and magnetization performance at room temperature do not improve due to the difficulty in aligning the magnetic moments, reflecting that the microstructure of the main phase 10 and sub-phase 20 is not optimal. In addition, since the magnetic properties depend on the microstructure of the diffusion precursor, diffusing Tb, a heavy rare earth element RH, into such a diffusion precursor does not improve the magnetic properties.

Comparative Example 13 is a sample of a rare-earth sintered magnet 1 with 0.15 at.% Tb diffused, manufactured according to a manufacturing method including hot working described in Patent Document 2, using Nd, La, Sm, Fe, and FeB as raw materials such as (Nd, La, Sm)-Fe-B. The composition ratio of Nd, La, and Sm is different from that of Comparative Example 12. When the microstructure of this sample is observed according to the method described above, the core-shell structure of the main phase 10 cannot be confirmed due to the absence of Pr addition. Furthermore, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, the second sub-phase 22 does not exist. Moreover, it cannot be confirmed that the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22. When the magnetic properties of this sample are evaluated according to the method described above, the coercivity and temperature coefficient of coercivity are "good" due to the refinement of the magnetic powder. Furthermore, because the magnetic moments are difficult to align, the residual magnetic flux density and magnetization performance are "poor". The temperature coefficient of the residual magnetic flux density is "good". This is because, although the presence of La and Sm in the main phase 10 or sub-phase 20 results in a good temperature coefficient for the magnetic properties, the residual magnetic flux density and magnetization performance at room temperature do not improve due to the difficulty in aligning the magnetic moments, reflecting that the microstructure of the main phase 10 and sub-phase 20 is not optimal. Even if the composition ratio of Nd, La, and Sm is changed, results almost the same as in Comparative Example 12 are obtained. Also, since the magnetic properties depend on the microstructure of the diffusion precursor which is the base material, diffusing Tb, a heavy rare earth element RH, into such a diffusion precursor does not improve the magnetic properties.

Comparative Example 14 is a sample of a rare earth sintered magnet 1 with 0.15 at.% Tb diffused, prepared according to a manufacturing method including hot working described in Patent Document 2, using Nd, Pr, La, Sm, Fe, and FeB as raw materials such as (Nd, Pr, La, Sm)-Fe-B. When the microstructure of this sample is observed according to the method described above, a core-shell structure is confirmed due to the addition of Pr and the hot working, but the core-shell structure is confirmed only in one type of main phase 10 where the Pr concentration in the core is high. Furthermore, although the concentration of Sm is segregated into one sub-phase 20 due to the segregation of La, a second sub-phase 22 does not exist. Moreover, it cannot be confirmed that the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22. Furthermore, when the magnetic properties of this sample are evaluated according to the method described above, the coercivity and temperature coefficient of coercivity are "good" due to the refinement of the magnetic powder. However, the residual magnetic flux density and magnetization performance are "poor" because the magnetic moments are difficult to align. The temperature coefficient of residual magnetic flux density is "good". This is because the formation of a core-shell structure with a high Pr concentration in the core significantly improves the coercivity to the level of rare-earth sintered magnet 1 with added Dy, and the presence of La and Sm in the main phase 10 or sub-phase 20 results in good temperature coefficients of magnetic properties, especially the temperature coefficient of coercivity. However, the residual magnetic flux density at room temperature does not improve because the magnetic moments are difficult to align, which reflects that the microstructure of the main phase 10 and sub-phase 20 is not optimal. Also, since the magnetic properties depend on the microstructure of the diffusion precursor, which is the base material, diffusing Tb, a heavy rare-earth element RH, into such a diffusion precursor does not improve the magnetic properties.

The samples of Examples 1 to 8 are rare earth sintered magnets 1 in which the heavy rare earth elements RH are Dy and Tb, R is one or more rare earth elements selected from those other than Nd, Pr, Dy, and Tb, satisfy the general formula (Nd, Pr, RH, R)-Fe-B, and have a main phase 10 containing crystal grains based on the Nd ₂ Fe ₁₄ B crystal structure, the main phase 10 has core portions 11c, 12c and shell portions 11s, 12s covering the core portions 11c, 12c, the main phase 10 is a mixture of a first main phase 11 where CNd > CPr and a second main phase 12 where CNd < CPr, and the concentration of the heavy rare earth element RH Dy in the core portion 11c of the first main phase 11 is higher than the concentration of Dy in the core portion 12c of the second main phase 12. Furthermore, the rare earth sintered magnets 1 of Examples 1 to 8 are characterized in that, when R = La, Sm, in addition to the first main phase 11 and the second main phase 12, they have a crystalline first sub-phase 21 based on an oxide phase whose main component is expressed as (Nd, Pr, Dy, Tb, La, Sm)-O, and a crystalline second sub-phase 22 whose main component is expressed as (Nd, Pr, Dy, Tb, La)-O, and the concentration of Sm is higher in the first sub-phase 21 than in the second sub-phase 22, forming an Sm-enriched portion 41 in the first sub-phase 21. When the magnetic properties of the samples of Examples 1 to 8 are evaluated according to the method described above, the remanent magnetic flux density is "good", the coercivity is "good", the temperature coefficient of the remanent magnetic flux density is "good", the temperature coefficient of the coercivity is "good", and the magnetization performance is "equal to or better". As a result, these rare earth sintered magnets 1 exhibit superior magnetic properties compared to conventional magnets, while reducing the use of Nd and heavy rare earth elements Dy and Tb, which are expensive, geographically unevenly distributed, and pose procurement risks. Furthermore, since magnetic properties depend on the microstructure of the diffusion precursor, diffusing Tb, a heavy rare earth element RH, into a diffusion precursor with good magnetic properties further improves the magnetic properties. In addition, in Examples 1 to 8, rare earth sintered magnets 1 with good magnetic properties can be obtained with a diffusion amount of 0.10 at.%, which is lower than the 0.15 at.% diffusion amount of Tb in Comparative Examples 1 to 14. In other words, compared to Comparative Examples 1 to 14, it is possible to obtain rare earth sintered magnets 1 that can significantly improve coercivity without reducing the residual magnetic flux density while reducing the amount of heavy rare earth element RH used.

Furthermore, as can be seen from Table 3, the samples of Examples 1 to 8 have a lower content of Dy, a heavy rare earth element RH in the general formula, and a lower diffusion amount of Tb, another heavy rare earth element RH, compared to the samples of Comparative Examples 1, 4, 8, and 11. However, the magnetic properties are superior to those of the samples of Comparative Examples 1, 4, 8, and 11. In particular, in order to obtain magnetic properties equivalent to those of the samples of Examples 1 to 8 with the samples of Comparative Examples 8 and 11, which are manufactured using the method described in Patent Document 2, it is necessary to include even more Dy and Tb. For this reason, the rare earth sintered magnets 1 according to Embodiments 1, 2, and 3 have the effect of reducing the amount of heavy rare earth elements RH used compared to the technology described in Patent Document 2. Alternatively, if the content of heavy rare earth element RH in the rare earth magnet manufactured by the manufacturing method of Patent Document 2 is the same as the content of heavy rare earth element RH in the rare earth sintered magnet 1 according to Embodiments 1, 2, and 3, the magnetic properties of the rare earth magnet manufactured by the manufacturing method of Patent Document 2 will be lower than those of the rare earth sintered magnet 1 according to Embodiments 1, 2, and 3. Furthermore, Table 3 shows that, compared to the samples of Comparative Examples 1 and 4 manufactured by the manufacturing method of Patent Document 1, the samples of Examples 1 to 8 can improve magnetic properties while reducing the content of Dy, which is a heavy rare earth element RH in the general formula, and also reducing the diffusion amount of Tb, which is a heavy rare earth element RH.

The configurations shown in the above embodiments are merely examples, and can be combined with other known technologies, or the embodiments themselves can be combined. Furthermore, parts of the configuration can be omitted or modified without departing from the spirit of the invention.

1 Rare earth sintered magnet, 10 Main phase, 11 First main phase, 11c, 12c Core, 11s, 12s Shell, 12 Second main phase, 20 Sub-phase, 21 First sub-phase, 22 Second sub-phase, 31 Heavy rare earth element-containing layer, 41 Sm-enriched section, 42 Heavy rare earth element-containing section, 100 Rotor, 101 Rotor core, 102 Magnet insertion hole, 120 Rotating machine, 130 Stator, 131 Teeth, 132 Winding.

Claims (10)

When RH is a heavy rare earth element containing at least one of Dy and Tb, and R is one or more rare earth elements selected from those other than Nd, Pr, Dy, and Tb, the main phase contains crystal grains that satisfy the general formula (Nd, Pr, RH, R)-Fe-B and are based on the Nd ₂ Fe ₁₄ B crystal structure,
Sub-phases existing between multiple main phases,
Equipped with,
The main phase has a core portion and a shell portion that covers the core portion.
The main phase comprises a first main phase where CNd > CPr and a second main phase where CNd < CPr, where the concentration of Nd in the core is CNd and the concentration of Pr in the core is CPr.
The concentration of heavy rare earth elements RH in the core portion of the first main phase is higher than the concentration of heavy rare earth elements RH in the core portion of the second main phase.
The first main phase and the second main phase are mixed together,
A rare earth sintered magnet characterized in that heavy rare earth elements are present on at least a portion of the surfaces of the first main phase and the second main phase. When RH is a heavy rare earth element containing at least one of Dy and Tb, and R is one or more rare earth elements selected from those other than Nd, Pr, Dy, and Tb, the main phase contains crystal grains that satisfy the general formula (Nd, Pr, RH, R)-Fe-B and are based on the Nd ₂ Fe ₁₄ B crystal structure,
Sub-phases existing between multiple main phases,
Equipped with,
The main phase has a core portion and a shell portion that covers the core portion.
The main phase comprises a first main phase where CNd > CPr and a second main phase where CNd < CPr, where the concentration of Nd in the core is CNd and the concentration of Pr in the core is CPr.
The concentration of heavy rare earth elements RH in the core portion of the first main phase is higher than the concentration of heavy rare earth elements RH in the core portion of the second main phase.
The first main phase and the second main phase are mixed together,
A rare earth sintered magnet characterized by the presence of heavy rare earth elements in the aforementioned subphase. The rare earth sintered magnet according to claim 1 or 2, characterized in that the number of first main phases is greater than the number of second main phases. When the concentration of Nd in the shell portion is SNd, the concentration of Pr in the shell portion is SPr, the concentration of heavy rare earth element RH in the core portion of the main phase is CRH, and the concentration of heavy rare earth element RH in the shell portion is SRH,
The rare earth sintered magnet according to claim 1 or 2, characterized in that the first main phase satisfies the relationship CNd > SNd, CPr < SPr, CRH > SRH, and the second main phase satisfies the relationship CNd < SNd, CPr > SPr, CRH < SRH. The aforementioned subphase comprises a crystalline first subphase based on an oxide phase whose main component is represented as (Nd, Pr, RH, La, Sm)-O when R is La, Sm, and a crystalline second subphase whose main component is represented as (Nd, Pr, RH, La)-O.
The rare earth sintered magnet according to claim 1 or 2 , characterized in that the concentration of Sm is higher in the first sub-phase than in the second sub-phase. The rare earth sintered magnet according to claim 5, characterized in that the first sub-phase has an Sm-enriched portion with a higher Sm concentration compared to the second sub-phase. The rare earth sintered magnet according to claim 6, characterized in that, in the first subphase, the heavy rare earth element is present so as to surround the outer periphery of the Sm-enriched portion. A method for manufacturing a rare earth sintered magnet according to claim 1 or 2 ,
A rare earth sintered magnet alloy manufacturing process for manufacturing a rare earth sintered magnet alloy that serves as a raw material for a diffusion precursor before diffusing the heavy rare earth elements into the rare earth sintered magnet,
A diffusion precursor manufacturing process for producing the aforementioned diffusion precursor,
A diffusion step of diffusing the heavy rare earth element into the diffusion precursor,
A cooling step of cooling the diffusion precursor in which the heavy rare earth element has been diffused,
Includes,
The aforementioned rare earth sintered magnet alloy manufacturing process is as follows:
A melting step of melting the raw materials for a rare earth sintered magnet alloy containing the elements constituting the diffusion precursor,
A first alloy cooling step is performed in which the molten raw material from the melting step is cooled to obtain a solidified alloy,
A second alloy cooling step is performed to further cool the solidified alloy to obtain a rare earth sintered magnet alloy,
Includes,
The aforementioned diffusion precursor manufacturing process is as follows:
A grinding step for grinding the rare earth sintered magnet alloy satisfying (Nd, Pr, RH, R)-Fe-B,
A molding step is performed to prepare a molded body by molding the powder of the rare earth sintered magnet alloy that has been pulverized in the pulverization step,
A sintering step to obtain a sintered body by sintering the molded body at a predetermined sintering temperature,
A first aging step in which the sintered body is held at a first aging temperature which is below the sintering temperature,
A second aging step is performed in which the sintered body held in the first aging step is held at a second aging temperature which is a temperature lower than the first aging temperature,
A third aging step in which the sintered body held in the second aging step is again held at the first aging temperature,
A fourth aging step in which the sintered body held in the third aging step is held at the second aging temperature,
A sintered body cooling step is performed to cool the sintered body held in the fourth aging step to obtain the diffusion precursor,
Includes,
A method for manufacturing a rare earth sintered magnet, characterized in that, in the diffusion step, the diffusion precursor is heat-treated at a temperature below the sintering temperature under conditions in which the diffusion precursor and the heavy rare earth element are present. Rotor core and
A rare earth sintered magnet according to claim 1 or 2 is provided in the rotor core,
A rotor characterized by having the following features. The rotor according to claim 9,
An annular stator is positioned opposite the rotor, having windings attached to teeth protruding toward the rotor on the inner surface of the side where the rotor is located,
A rotating machine characterized by having the following features.