Anisotropic rare earth sintered magnet and method for producing the same

Abstract

To provide an anisotropic rare earth sintered magnet indicating excellent magnetic characteristics in the anisotropic rare earth sintered magnet defining a compound of an Nd2Fe14B type crystal as a main phase and containing Ce, and a manufacturing method thereof.SOLUTION: The present invention relates to an anisotropic rare earth sintered magnet and a composition of the anisotropic rare earth sintered magnet is expressed by a formula Rx(Fe1-aCoa)100-x-y-zByMz (R is two or more kinds of elements selected from rare earth elements and essentially containing Nd and Ce). In the anisotropic rare earth sintered magnet, a main phase consists of a compound of an Nd2Fe14B type crystal. A main phase particle 10 exists in which a Ce / R' ratio (R' is one or more kinds of elements selected from rare earth elements and essentially containing Nd) in a central part of the particle is lower than a Ce / R' ratio in an outer shell part of the particle, and an R' rich phase and an R'(Fe, Co)2 phase containing Ce exist in a particle boundary part 20. The present invention also relates to a manufacturing method of the anisotropic rare earth sintered magnet.SELECTED DRAWING: Figure 1

Description

Technical Field

[0001] The present invention relates to an anisotropic rare earth sintered magnet having a compound of Nd 2 Fe 14 B type crystal as a main phase and containing Ce, and a method for manufacturing the same.

Background Art

[0002] Nd-Fe-B sintered magnets are expected to have an increasing demand and production volume in the future, against the backdrop of the electrification of automobiles and the high performance and power saving of industrial motors. However, rare earth elements such as Nd, Pr, Dy, and Tb used as raw materials are expensive and also have risks to future supply stability. For this reason, research has been conducted to replace part of Nd with Ce, which has a higher elemental content rate in the earth's crust and is inexpensive.

[0003] For example, in Patent Document 1, it includes a main phase and a grain boundary phase, and the overall composition is (R 2 (1-x) R 1 x ) y Fe (100-y-w-z-v) Co w B z M 1 v ·(R 3 (1-p) M 2 p ) q (wherein, R 1 is an element selected from Ce, La, Y, and Sc, R 2 and R 3 are elements selected from Nd, Pr, Gd, Tb, Dy, and Ho, M 1 is a predetermined element, etc., M 2 is a transition metal element alloyed with R 3 ), and the main phase is R 2 Fe 14It has a B-type crystal structure, the average particle size of the main phase is 1 to 20 μm, the main phase has a core part and a shell part, the thickness of the shell part is 25 to 150 nm, and when the light rare earth element ratio of the core part is a and the light rare earth element ratio of the shell part is b, a rare earth magnet excellent in both coercive force and residual magnetization that satisfies 0 ≦ b ≦ 0.30 and 0 ≦ b / a ≦ 0.50 and a method for manufacturing the same are shown.

[0004] In Patent Document 2, a main phase particle containing R, T, and B and a grain boundary phase are provided. R contains Nd and Ce, T contains Fe, the grain boundary phase contains an R-T phase and an R-rich phase, the R-T phase contains an intermetallic compound of R and T, the content of R in the R-rich phase is larger than the content of R in the R-T phase, Ce / R×100 = 65 to 100 in the R-T phase, and the content of R in the R-rich phase is 70 to 100 atomic %. A rare earth magnet is shown.

[0005] In Patent Document 3, the overall composition is of the formula (Nd (1-x-y) Ce x R 1 y ) p (Fe (1-z) Co z ) (100-p-q-r-s) B q Ga r M s (where R 1 is one or more selected from rare earth elements other than Nd and Ce and Y, M is one or more selected from Al, Cu, Au, Ag, Zn, In, Mn, Zr, and Ti and inevitable impurity elements, and 12 ≦ p ≦ 20, 4.0 ≦ q ≦ 6.5, 0 ≦ r ≦ 1.0, 0 ≦ s ≦ 0.5, 0 < x ≦ 0.35, 0 ≦ y ≦ 0.10, and 0.050 ≦ z 0.140), and a rare earth magnet and a method for manufacturing the same that include a magnetic phase and a grain boundary phase existing around the magnetic phase are shown.

[0006] In Patent Document 4, a permanent magnet is provided which includes a plurality of main phase particles containing rare earth element R, transition metal element T, and boron B, and a grain boundary phase located between the plurality of main phase particles. R includes Nd and Ce, T includes Fe, the total content of R in the permanent magnet is [R] atomic %, the total content of T in the permanent magnet is [T] atomic %, the content of B in the permanent magnet is [B] atomic %, the content of Ce in the permanent magnet is [Ce] atomic %, [Ce] / [R] is 0.1 to 0.6, [T] / [B] is 14 to 18, the grain boundary phase includes an R-T phase containing an intermetallic compound of R and T, the area of the unit cross-section of the permanent magnet is A0, the total area of the R-T phase in the unit cross-section is AL, and AL / A0 is 0.05 to 0.5, showing a permanent magnet with high flexural strength.

[0007] In Patent Document 5, (Ce x Nd (1-x) ) y Fe (100-y-w-z-v) Co w B z M v (where M is at least one of Ga, Al, Cu, Au, Ag, Zn, In, Mn, 0 ≦ x ≦ 0.75, 5 ≦ y ≦ 20, 4 ≦ z ≦ 6.5, 0 ≦ w ≦ 8, 0 ≦ v ≦ 2), showing a rare earth magnet including crystal grains having the overall composition, which are composed of a core part 1 and a shell part 2 around it, and having a higher Nd concentration in the shell part 2 than in the core part 1.

[0008] In Patent Document 6, in an R-T-B based magnet which requires R1 and Ce as R, by subjecting the raw material R-T-B based magnet to heat treatment for a long time, the main phase particles are core-shelled. When the mass concentrations of R1 and Ce in the core part are αNd and αCe respectively, and the mass concentrations of R1 and Ce in the shell part are βR1 and βCe respectively, a ratio (βR1 / βCe = B) of the mass concentrations of R1 and Ce in the shell part to a ratio (αR1 / αCe = A) of the mass concentrations of R1 and Ce in the core part (B / A) is 1.1 or more, showing an R-T-B based sintered magnet which improves the adhesion strength with plating while suppressing the decrease in coercivity by adding Ce.

Prior Art Documents

Patent Documents

[0009]

Patent Document 1

Patent Document 2

Patent Document 3

Patent Document 4

Patent Document 5

Patent Document 6

Summary of the Invention

Problems to be Solved by the Invention

[0010] As described above, in the R-T-B-based magnet containing Ce, it has been proposed that good properties can be obtained by providing main phase grains having a core-shell structure or using an R-T intermetallic compound as the grain boundary phase. However, for the R 2 Fe 14 magnetic properties of the B compound at room temperature, when R = Nd, the saturation magnetization M s is 1.60 T and the anisotropy magnetic field μ 0 H A is 6.7 T, while when R = Ce, M s is 1.17 T and μ 0 H A is 3.0 T, which is low. Therefore, it is difficult to solve the problem that the magnet properties deteriorate when the Ce content is high.

[0011] The present invention has been made in view of the above problems, and an object of the present invention is to provide an anisotropic rare earth sintered magnet showing good magnetic properties and a method for manufacturing the same in an anisotropic rare earth sintered magnet having a compound of Nd 2 Fe 14 B type crystal as the main phase and containing Ce.

Means for Solving the Problems

[0012] As a result of repeated studies to achieve the above object, the inventors of the present invention have found that in an anisotropic rare earth sintered magnet having a compound of Nd 2 Fe 14 B type crystal as the main phase and containing Ce, there are main phase grains in which the Ce / R' ratio (R' is one or more elements selected from rare earth elements and including Nd as an essential element) in the central part of the grains is lower than the Ce / R' ratio in the outer shell part of the grains, and an R' rich phase containing Ce and an R'(Fe,Co) 2 phase containing Ce are present at the grain boundary, and good magnetic properties can be obtained. Thus, the present invention has been completed.

[0013] Therefore, the present invention provides the following anisotropic rare earth sintered magnet and a method for manufacturing the same. (1) An anisotropic rare earth sintered magnet having a composition represented by the formula R x (Fe 1-a Co a ) 100-x-y-z B y M z (where R is two or more elements selected from rare earth elements and including Nd and Ce as essential elements, M is one or more elements selected from the group consisting of Al, Si, Ti, V, Cr, Mn, Ni, Cu, Zn, Ga, Ge, Zr, Nb, Mo, Ag, In, Sn, Hf, Ta, W, Pb, Bi, and x, y, z, a are 12 ≦ x ≦ 17 atomic %, 3.5 ≦ y ≦ 6.0 atomic %, 0 ≦ z ≦ 3 atomic %, 0 ≦ a ≦ 0.1, respectively), and the main phase is composed of a compound of Nd 2 Fe 14 B type crystal, there are main phase grains in which the Ce / R' ratio (R' is one or more elements selected from rare earth elements and including Nd as an essential element) in the central part of the grains is lower than the Ce / R' ratio in the outer shell part of the grains, and an R' rich phase containing Ce and an R'(Fe,Co) 2 phase are present at the grain boundary. (2) The main phase and the R'(Fe,Co) 2The anisotropic rare earth sintered magnet according to (1), wherein a boundary phase containing 20 atomic % or more of R and having a thickness of 20 nm or less is formed between the phases. (3) The anisotropic rare earth sintered magnet according to (1) or (2), wherein in the main phase grains, there are main phase grains in which the central R' does not contain Ce. (4) The anisotropic rare earth sintered magnet according to any one of (1) to (3), wherein in the main phase grains, there are main phase grains in which the central R' is Nd or consists of Nd and Pr. (5) The R'(Fe,Co) 2 phase is a phase showing ferromagnetic or ferrimagnetic properties at room temperature or higher, and the anisotropic rare earth sintered magnet according to any one of (1) to (4). (6) The R'(Fe,Co) 2 The anisotropic rare earth sintered magnet according to any one of (1) to (5), wherein the Ce / R' ratio in the phase is higher than the Ce / R' ratio in the outer shell part of the main phase grains. (7) The anisotropic rare earth sintered magnet according to any one of (1) to (6), wherein the Ce / R' ratio in the R'-rich phase is higher than the Ce / R' ratio in the outer shell part of the main phase grains. (8) The anisotropic rare earth sintered magnet according to any one of (1) to (7), wherein the R'-rich phase and the R'(Fe,Co) 2 phase are contained in a total amount of 1 volume % or more. (9) The anisotropic rare earth sintered magnet according to any one of (1) to (8), wherein the Ce / R' ratio in the composition of the sintered body is 0.01 or more and 0.3 or less. (10) The anisotropic rare earth sintered magnet according to any one of (1) to (9), wherein the B-rich phase contained in the sintered magnet is 5 volume % or less. (11) The anisotropic rare earth sintered magnet according to any one of (1) to (10), wherein an intergranular phase between two particles is formed between adjacent main phase grains. (12) The main phase and the R'(Fe,Co) 2The anisotropic rare-earth sintered magnet according to (11), wherein Ce / R’ in the boundary phase formed between the phases is higher than Ce / R’ in the intergranular phase between two adjacent main-phase grains. (13) Coercive force H at room temperature cJ(room) is 10 kOe or more, and the value of the temperature coefficient β of the coercive force is β ≧ (0.01 × H cJ(室温) - 0.720)% / K, which is the anisotropic rare-earth sintered magnet according to any one of (1) to (12). (14) Nd 2 Fe 14 An alloy containing a compound phase of NdFeB-type crystal, an alloy having a higher R’ composition ratio and Ce / R’ ratio than that, are pulverized and mixed, and compacted under a magnetic field application to form a compact, and then sintered at a temperature of 800 °C or more and 1200 °C or less, which is the method for manufacturing an anisotropic rare-earth sintered magnet according to any one of (1) to (13). (15) Nd 2 Fe 14 An alloy containing a compound phase of NdFeB-type crystal is pulverized, compacted under a magnetic field application to form a compact, and then sintered at a temperature of 800 °C or more and 1200 °C or less. A material containing Ce is brought into contact with the sintered body, and heat treatment is performed at a temperature of 600 °C or more and the sintering temperature or less to diffuse Ce into the interior of the sintered body, which is the method for manufacturing an anisotropic rare-earth sintered magnet according to any one of (1) to (14). (16) The material containing Ce brought into contact with the sintered body is one or more selected from Ce metal, Ce-containing alloy, and Ce-containing compound, and the form thereof is one or more selected from powder, thin film, thin strip, foil, and gas, which is the method for manufacturing an anisotropic rare-earth sintered magnet according to (15). (17) The method for manufacturing an anisotropic rare-earth sintered magnet according to any one of (14) to (16), wherein heat treatment is performed on the sintered body at a temperature of 300 to 800 °C. (18) After heat treatment is performed on the sintered body at a temperature of 600 to 1000 °C, it is cooled at a cooling rate of 1 °C / min or more and 50 °C / min or less to at least 550 °C or less, and further heat treatment is performed at a temperature of 300 to 800 °C, which is the method for manufacturing an anisotropic rare-earth sintered magnet according to any one of (14) to (17).

Advantages of the Invention

[0014] According to the present invention, in an anisotropic rare earth sintered magnet having a compound of Nd 2 Fe 14 B type crystal as the main phase and containing Ce, an anisotropic rare earth sintered magnet showing good magnetic properties can be obtained.

Brief Description of the Drawings

[0015]

Figure 1

Figure 2

Figure 3

Figure 4

Modes for Carrying Out the Invention

[0016] Hereinafter, embodiments of the present invention will be described. The magnet of the present invention has a composition represented by the following formula R x (Fe 1-a Co a ) 100-x-y-z B y M z and a compound of Nd 2 Fe 14 B type crystal is the main phase, and particles having different Ce / R' ratios exist in the central part and the outer shell part of the grains in the main phase particles, and an R'-rich phase containing Ce and an R'(Fe, Co) containing Ce are present at the grain boundary 2It is an anisotropic rare earth sintered magnet having a phase. First, each component will be described below. Here, R is two or more elements selected from rare earth elements and including Nd and Ce as essential elements, and M is one or more elements selected from the group consisting of Al, Si, Ti, V, Cr, Mn, Ni, Cu, Zn, Ga, Ge, Zr, Nb, Mo, Ag, In, Sn, Hf, Ta, W, Pb, and Bi. Also, x, y, z, and a are 12 ≦ x ≦ 17 atomic %, 3.5 ≦ y ≦ 6.0 atomic %, 0 ≦ z ≦ 3 atomic %, and 0 ≦ a ≦ 0.1, respectively. Further, R' is one or more elements selected from rare earth elements and including Nd as an essential element. Note that the R'-rich phase is a phase in which R' is contained in an amount exceeding 40 atomic %. Also, R'(Fe,Co) 2 phase is MgCu 2 It has a structure and is a compound phase called a Laves phase.

[0017] As described above, R is two or more elements selected from rare earth elements and including Nd and Ce as essential elements. Specifically, R necessarily contains Nd and Ce, and may further contain one or more elements selected from Sc, Y, La, Pr, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu. R is an element necessary for forming a compound having a Nd 2 Fe 14 B-type crystal structure. The content of R is set to be 12 atomic % or more and 17 atomic % or less. If it is 12.5 atomic % or more and 16 atomic % or less, it is more preferable. If it is less than 12 atomic %, the α-Fe phase precipitates and sintering becomes difficult. On the other hand, if it exceeds 17 atomic %, the volume ratio of the Nd 2 Fe 14 B-type compound phase decreases and good magnetic properties cannot be obtained. Nd 2 Fe 14Since B-type compounds exhibit particularly high magnetic properties when R is Nd, the anisotropic rare earth sintered magnet of the present invention requires Nd. Furthermore, in order to reduce the cost of the magnet and stabilize the supply of elements, Ce, which has a high elemental abundance ratio among rare earth elements, must be included. The amount of Ce contained in R in the sintered body composition is preferably 1% to 30% of R in atomic ratio, more preferably 3% to 25%, and particularly preferably 5% to 20%. By having the Ce ratio in this range, a high residual magnetic flux density B r and high coercive force H cJ , and even better H cJ An anisotropic sintered magnet with excellent temperature characteristics can be obtained.

[0018] B and Nd 2 Fe 14 It is an element essential for forming a B type compound. The content of B is 3.5 atomic % or more and 6.0 atomic % or less, and more preferably 5.0 atomic % or more and 5.8 atomic % or less. If it is less than 3.5 atomic %, R 2 Fe 17 On the other hand, if it exceeds 6.0 atomic %, a different phase such as a B-rich phase is formed, the volume ratio of the main phase decreases, and good magnetic properties cannot be obtained.

[0019] As mentioned above, M is one or more elements selected from Al, Si, Ti, V, Cr, Mn, Ni, Cu, Zn, Ga, Ge, Zr, Nb, Mo, Ag, In, Sn, Hf, Ta, W, Pb and Bi. These elements include Nd 2 Fe 14 The H cJ However, if contained in excess, it will lower the Br of the magnet. Therefore, when M is contained, its total content is set to 3 atomic % or less. 2 atomic % or less is more preferable, and 1 atomic % or less is particularly preferable.

[0020] The anisotropic rare earth sintered magnet of the present invention contains Fe as an essential component along with R and B. Furthermore, a part of Fe may be substituted with Co. The substitution with Co is carried out by replacing Nd 2 Fe14 has the effect of increasing the Curie temperature T of the B-type compound. The substitution rate of Co is 10% or less in atomic ratio. When the substitution rate exceeds 10%, M c will conversely decrease. The proportions of Fe and Co are the remainder of R, B, and M. In addition to this, it may contain inevitable impurities incorporated from raw materials or mixed in the manufacturing process, specifically H, C, N, O, F, P, S, Mg, Cl, Ca, etc. From the viewpoint of obtaining good magnetic properties, the total content is preferably 3% by weight or less, more preferably 1% by weight or less. In particular, C, N, and O preferably total 1% by weight or less, more preferably 0.5% by weight or less, and particularly preferably 0.3% by weight or less. s

[0021] Next, the phases constituting the anisotropic rare earth sintered magnet of the present invention will be described. The main phase in the anisotropic rare earth sintered magnet of the present invention is Nd 2 Fe 14 and consists of a compound having a B-type crystal structure. The average crystal grain size of the main phase is preferably 1 μm or more and 15 μm or less. A range of 1.5 μm or more and 10 μm or less is more preferable, and a range of 2 μm or more and 5 μm or less is particularly preferable. By setting the average crystal grain size within such a range, a decrease in the residual magnetic flux density B r due to a decrease in the degree of orientation of crystal grains and a decrease in the coercive force H cJ can be suppressed. From the viewpoint of obtaining good B r and H cJ , the volume ratio of the main phase is preferably 80% by volume or more and less than 99% by volume with respect to the entire magnet, and more preferably 90% by volume or more and 99% by volume or less. Regarding the crystal grain size of the main phase, the cross-section of the sintered magnet is polished until it becomes mirror-like, immersed in an etching solution (such as a mixed solution of nitric acid + hydrochloric acid + glycerin) to selectively remove the grain boundary phase, and then observed with a laser microscope at any 10 or more locations on this cross-section. The cross-sectional area of each particle is calculated by image analysis from the obtained observation image, and the average diameter when these are regarded as circles is calculated as the average crystal grain size. ​Also, regarding the volume ratios of the main phase and each phase, after polishing the cross-section of the sintered magnet until it becomes mirror-like, EPMA is used to observe the microstructure of the anisotropic rare-earth sintered magnet and analyze the composition of each phase. After confirming the presence of the main phase, R'-rich phase, and R'(Fe,Co) 2 phase, it can be calculated that the area ratio in the image of the backscattered electron image is equal to the volume ratio of each phase.

[0022] R’ 2 Fe 14 The B compound has the highest saturation magnetization M when R' = Nd. When a part of Nd is substituted with Ce, M s decreases as the Ce substitution amount increases. Therefore, in the magnet of the present invention, in order to reduce the influence of the decrease in B of the magnet due to Ce substitution, the Ce / R' ratio (atomic ratio of Ce to R') is different between the central part and the outer shell part of the main phase grains, and there are main phase grains in which the Ce / R' ratio in the central part of the grains is lower than the Ce / R' ratio in the outer shell part of the grains. However, main phase grains with a uniform Ce concentration distribution may be partially included. Here, the outer shell part refers to the region including the surface of the main phase grains, and the central part refers to the other internal region. By taking such a microstructure, the decrease in M s is suppressed in the region near the center of the main phase grains with a low Ce / R' ratio, and the amount of decrease in B r of the magnet due to Ce substitution can be reduced. It is more preferable when Ce is not contained in R' in the central part of the main phase grains, and even more preferable when R' in the grain central part consists of Nd or Nd and Pr. s r On the other hand, as will be described later, when an R'-rich phase containing Ce and an R'(Fe,Co) 2 phase are formed at the grain boundary, H cJ increases at room temperature, and the temperature change of H cJ becomes smaller, showing excellent magnetic properties. In order to efficiently form these phases, in the magnet of the present invention, the structure is such that the Ce / R' ratio of the outer shell part of the main phase grains is higher than the Ce / R' ratio of the central part of the main phase grains. As a result, the Ce concentration at the grain boundary also increases, and the R'(Fe,Co) 2 phase is likely to be formed at the grain boundary. On the contrary, when the Ce / R' ratio of the grains is uniform, R'(Fe,Co) 2 phase is less likely to be formed at the grain boundary. When the Ce / R' ratio of the grains is uniform, R'(Fe,Co)phase is less likely to be formed at the grain boundary.2 In order to significantly form a phase, it is necessary to increase the Ce substitution amount of the sintered body, which causes a significant decrease in M s . When the Ce / R' ratio in the outer shell part of the grains is high, H on the grain surface A decreases, but due to the large H-increasing effect of the R'-rich phase containing Ce and the R'(Fe,Co) 2 phase, the negative effect caused by the decrease in H cJ is reduced. A Conversely, when there are main phase grains in which the Ce / R' ratio in the central part of the grains is higher than the Ce / R' ratio in the outer shell part of the grains, M in the region near the center of the main phase grains with a high Ce / R' ratio s decreases significantly, which is not compatible with the guidelines for the magnets of the present invention. Therefore, in the magnets of the present invention, it is assumed that there are no main phase grains in which the Ce / R' ratio in the central part of the grains is higher than the Ce / R' ratio in the outer shell part of the grains. Although the thickness of the outer shell part with a high Ce / R' ratio is not particularly limited, from the viewpoint of increasing the volume ratio of the inner part of the outer shell part, 1 nm to 2 μm is preferable, and 2 nm to 1 μm is particularly preferable.

[0023] The R'-rich phase and R'(Fe,Co) 2 phase are formed at the grain boundaries of the magnet structure. The grain boundaries include, in addition to the grain boundary triple points, grain boundary phases between two particles, etc. Here, it is assumed that the phase contains more than 40 atomic% of R'. The inventors have found that when the R'-rich phase containing Ce and the R'(Fe,Co) 2 phase exist at the grain boundaries, H at room temperature of the magnet cJ is improved, and furthermore, the temperature characteristics of H cJ are also improved. In order to obtain a structure in which the two phases coexist, the Ce / R' ratio in the composition of the sintered body is preferably 0.01 or more and 0.3 or less. If it is less than 0.01, the R'(Fe,Co) 2 phase is not formed, and if it exceeds 0.3, it becomes difficult for the R'-rich phase to exist. More preferably, it is 0.03 or more and 0.25 or less, and particularly preferably 0.05 or more and 0.2 or less.

[0024] The R'-rich phase and R'(Fe,Co)2 The phase mainly brings about four effects. The first effect is the action of promoting sintering. At the sintering temperature, both the R'-rich phase and the R'(Fe,Co) 2 phase also melt to form a liquid phase. Therefore, liquid-phase sintering proceeds, and sintering is completed more rapidly compared to solid-phase sintering when these phases are not included. Also, due to the coexistence of the R'-rich phase and the R'(Fe,Co) 2 phase, the liquid-phase formation temperature tends to drop compared to the case of only one of the phases, and liquid-phase sintering proceeds more rapidly.

[0025] The second effect is the cleaning of the main-phase grain surface. Since the anisotropic rare-earth sintered magnet of the present invention has a nucleation-type coercivity mechanism, it is desirable that the surface of the main-phase grains be smooth so that nucleation of reverse magnetic domains is less likely to occur. The R'-rich phase and the R'(Fe,Co) 2 phase play a role in smoothing the surface of the main-phase grains during the sintering process or the subsequent aging process. By this cleaning effect, nucleation of reverse magnetic domains, which is a factor in reducing coercivity, is suppressed. The R'(Fe,Co) 2 phase has relatively high wettability with respect to the main phase compared to other compound phases such as R'M 3 、R'M 2 、R'(Fe,Co)M and R'(Fe,Co) 2 M 2 etc. In particular, this phase is more likely to coat the surface of the main-phase grains by coexisting with the R'-rich phase, resulting in a large cleaning effect. Thereby, nucleation of reverse magnetic domains is suppressed, the coercivity at room temperature increases, and the decrease in coercivity at high temperatures also becomes small, and it is considered that a good temperature dependence of H cJ is exhibited.

[0026] The third effect is the effect of weakening the magnetic interaction between the main-phase grains. The R'-rich phase and the R'(Fe,Co) 2In a magnet having a phase, by performing an optimum sintering treatment or aging treatment, a grain boundary phase containing more R' than the main phase is formed between adjacent main phase grains. As a result, the magnetic interaction between the main phase grains is weakened and the coercive force is manifested. When the grain boundary phase between two particles contains Ce, the effect of weakening the magnetic interaction between the main phase grains becomes greater, and it is considered to act in the direction of further increasing the coercive force.

[0027] The fourth effect is the effect of promoting the formation of a boundary phase between the R'(Fe,Co) 2 phase and the main phase. In a magnet in which an R'-rich phase and an R'(Fe,Co) 2 phase are present at the grain boundary, by optimally performing sintering and subsequent heat treatment according to conditions such as composition and powder particle size, not only between the main phase grains but also between the R'(Fe,Co) 2 phase and the main phase grains, a boundary phase with a thin thickness is formed. The R'(Fe,Co) 2 phase in the magnet of the present invention is a magnetic phase, but the formation of this thin boundary phase weakens the magnetic interaction between the R'(Fe,Co) 2 phase and the main phase, and a high coercive force can be obtained. In a magnet in which no R'-rich phase exists at the grain boundary, it is difficult to form a thin boundary phase between the R'(Fe,Co) 2 phase and the main phase grains or a grain boundary phase between two main phase grains, or it is difficult to form a structure in which the surface of the main phase grains is completely covered with these, so it is difficult to obtain a magnet showing sufficient coercive force.

[0028] As described above, the R'-rich phase shall contain at least more than 40 atomic% of R'. When R' exceeds 40 atomic%, the wettability with the main phase becomes better, and the above-mentioned effects are more easily obtained. It is more preferable to contain 50 atoms or more of R', and particularly preferable to contain 60 atoms or more of R'. The R'-rich phase may be an R' metal phase, an amorphous phase, or R' 3 (Fe,Co,M), R' 2 (Fe,Co,M), R' 5 (Fe,Co,M) 3, it may be an intermetallic compound with a high R' composition and a low melting point such as R’(Fe, Co, M). Also, the total content of Fe, Co, M elements and impurity elements such as H, B, C, N, O, F, P, S, Mg, Cl, Ca, etc. may be less than 60 atomic %. In addition, the higher the Ce / R ratio of the R-rich phase, the greater the effect of reducing the magnetic interaction between the main phase grains. Therefore, in order to make Ce effectively act on improving the magnetic properties, the Ce / R ratio in the R-rich phase is preferably higher than the Ce / R ratio in the outer shell part of the main phase grains.

[0029] On the other hand, R’(Fe, Co) 2 phase is a Laves compound of the MgCu 2 type crystal. However, when performing composition analysis using EPMA etc., considering measurement variations etc., it is assumed to contain R’ of 20 atomic % or more and less than 40 atomic %. Also, part of Fe and Co may be substituted by the M element. However, the substitution amount of M is within the range where the MgCu 2 type crystal structure is maintained.

[0030] The R’(Fe, Co) 2 phase in the anisotropic rare earth sintered magnet of the present invention is a magnetic phase. The magnetic phase referred to here shows ferromagnetic or ferrimagnetic, and is a phase where the Curie temperature T c is above room temperature (23 °C). R’Fe 2 except for CeFe 2 has T c above room temperature, and if 10% or more of CeFe 2 is substituted by other elements, T c will be above room temperature. On the other hand, R’Co 2 except for GdCo 2 has T c below room temperature or is a paramagnetic phase, but in the anisotropic rare earth sintered magnet of the present invention, since the substitution atomic ratio of Fe by Co is 0.1 or less, in most cases the R’(Fe, Co) 2 phase becomes a magnetic phase. Generally, the soft magnetic phase contained in the structure often has an adverse effect on the magnetic properties, but in the anisotropic rare earth sintered magnet of the present invention, R’(Fe, Co) 2The cleaning effect on the surface of the main phase grains due to the phase and the effect of forming the grain boundary phase between two particles are greater, and even for the magnetic phase, at room temperature H cJ increase and H cJ temperature dependence improvement are considered to contribute. Also, the R’(Fe,Co) 2 phase is difficult to exist stably when R’ is only Nd or Pr, and is formed at the grain boundary as an equilibrium phase by including Ce. Therefore, the Ce / R’ ratio of the R’(Fe,Co) 2 phase is preferably higher than the Ce / R’ ratio of the outer shell part of the main phase grains.

[0031] The formation amounts of the R’-rich phase and the R’(Fe,Co) 2 phase are preferably 1% by volume or more in total, more preferably 1% by volume or more and less than 20% by volume. Further preferably, it is 1.5% by volume or more and less than 15% by volume, and even more preferably in the range of 2% by volume or more and less than 10% by volume. Also, the R’-rich phase and the R’(Fe,Co) 2 phase are each preferably 0.5% by volume or more. By setting it in such a range, the area in contact with the main phase grains is ensured, and the effect of H cJ increase is easily obtained. Also, the decrease in B r is suppressed, and the desired magnetic properties are easily obtained.

[0032] In a more preferable structure of the sintered magnet of the present invention, a boundary phase having a thin thickness is formed between the R’(Fe,Co) 2 phase and the main phase. Since the R’(Fe,Co) 2 phase and the main phase are separated by this thin boundary phase, the magnetic interaction between the two phases is weakened, and the temperature dependence of room temperature H cJ and H cJ is further improved. This boundary phase may be in an amorphous state with a disordered atomic arrangement or may have regularity in the atomic arrangement. When observing the boundary phase using an apparatus such as STEM (scanning transmission electron microscope), its composition shall contain 20 atomic% or more of R'. If the content of R' is 20 atomic% or more, the effect of improving the coercive force due to the boundary phase is easily obtained. If the content of R' is 25 atomic% or more, it is more preferable, and if it is 30 atomic% or more, it is even more preferable. In addition, elements such as C, N, and O may be contained in addition to R', Fe, Co, and M. The thickness of the boundary phase is preferably 0.1 nm or more and 20 nm or less. Within such a range, for R(Fe, Co) 2 an effect occurs that weakens the magnetic interaction between the phase and the main phase, and it is also possible to suppress the decrease in the volume fraction of the main phase due to the formation of the boundary phase. A thickness of 0.2 nm or more and 10 nm or less is more preferable, and a thickness of 0.5 nm or more and 5 nm or less is particularly preferable.

[0033] R’(Fe, Co) 2 The Ce / R’ of this thin boundary phase formed between the R’(Fe, Co) phase and the main phase is preferably higher than the Ce / R’ of the two-particle intergranular phase formed between the main phase grains. Since the boundary phase is adjacent to the R’(Fe, Co) phase containing a large amount of Ce, it is easy to stably realize a high Ce / R’ composition. The higher the Ce / R’, the greater the effect of weakening the magnetic interaction. Therefore, when the area of the main phase grain surface covered by this phase increases, the magnet shows a higher room temperature H 2 as shown. The value of Ce / R’ of the boundary phase is preferably 0.2 or more. 0.3 or more is more preferable, and 0.35 or more is particularly preferable. cJ Thus, by taking an organizational form in which a boundary phase with a high Ce / R’ is formed between the main phase grains and the R’(Fe, Co) phase, the magnetic interaction between the main phase - R’(Fe, Co) phases is weakened, and a high room temperature H 2 and good H 2 temperature dependence are obtained. cJ and good H cJ temperature dependence are obtained.

[0034] Note that the above-mentioned R’(Fe, Co) 2The thickness of the boundary phase formed between the phase and the main phase, and the two-particle grain boundary phase between the main phase grains, for example, can be calculated from the obtained HAADF image (High-Angle Annular Dark Field) by observing the locations where the main phase grains are adjacent to each other and the locations where the R’(Fe,Co) 2 phase and the main phase are adjacent, using a STEM device (JEM-ARM200F manufactured by JEOL Ltd.).

[0035] In addition, the anisotropic rare earth sintered magnet of the present invention may contain R’ oxides, R’ carbides, R’ nitrides, R’ oxycarbides, M carbides, etc. formed by unavoidably mixed C, N, O. From the viewpoint of suppressing deterioration of magnetic properties, the volume ratio of these is preferably 10% by volume or less, and more preferably 5% by volume or less.

[0036] It is preferable that the phases other than the above are as few as possible. For example, the R’ 1+ε (Fe,Co) 4 B 4 rich phase represented by is preferably 5% by volume or less in order to suppress a decrease in the volume ratio of the main phase, the R’ rich phase, and the R’(Fe,Co) 2 phase. Also, from the viewpoint of preventing a significant decrease in magnetic properties, it is preferable that the α-(Fe,Co) phase and the R’ 2 (Fe,Co,M) 17 phase are not contained in the anisotropic rare earth sintered magnet of the present invention.

[0037] Next, the manufacturing method will be described. The anisotropic rare earth sintered magnet of the present invention is manufactured by powder metallurgy. Examples of means for manufacturing a magnet having a structure with different Ce / R’ ratios in the central part and the outer shell part of the main phase grains include, for example, the two-alloy method and the grain boundary diffusion method.

[0038] First, in order to produce the raw material alloy, metal raw materials such as R’, Fe, Co, and M, alloys, ferroalloys, etc. are used. Considering raw material losses during the manufacturing process, etc., the final sintered body is adjusted so as to have a predetermined composition. These raw materials are melted in a high-frequency furnace or an arc furnace, etc. to produce an alloy. Cooling from the molten metal may be by a casting method or as a thin sheet by a strip casting method. In the case of the strip casting method, it is preferable to adjust the cooling rate to produce an alloy such that the average crystal grain size of the main phase or the average grain boundary phase interval is 1 μm or more. If it is less than 1 μm, the powder after fine pulverization becomes polycrystalline, and the main phase crystal grains do not sufficiently orient in the magnetic field forming process, leading to a decrease in B r of. The average crystal grain size can be calculated, for example, by polishing the cross-section of the alloy, performing etching treatment, observing the microstructure, drawing 20 lines at equal intervals parallel to the roll contact surface, and counting the intersection points where these lines intersect the grain boundary phase part removed by etching. When α-Fe precipitates in the alloy, heat treatment may be applied to the alloy so that the formation amount of the Nd 2 Fe 14 B-type compound phase increases.

[0039] The above raw material alloy is roughly pulverized into powder with an average particle size of 0.05 - 3 mm by means such as mechanical pulverization with a brown mill or hydrogenation pulverization. Alternatively, the HDDR method (hydrogen disproportionation desorption recombination method) may be applied. Further, the coarse powder is finely pulverized by a ball mill or a jet mill using high-pressure nitrogen, etc. to obtain powder with an average particle size of 0.5 - 20 μm, more preferably 1 - 10 μm. Note that lubricants, etc. may be added before and after the fine pulverization process as necessary.

[0040] When using the two-alloy method, two types of raw material alloys with different compositions are produced. Note that three or more types of alloys may also be used. At this time, Nd 2 Fe 14It is preferable to combine alloy A with a relatively low Ce / R’ ratio mainly composed of a B-type compound phase and alloy B with a relatively higher R’ composition ratio and Ce / R’ ratio than alloy A, and adjust the average composition to a predetermined composition. These alloys are produced by a casting method or a strip casting method and then pulverized. The step of mixing the respective alloy powders may be carried out in the coarse powder state before fine pulverization or after fine pulverization.

[0041] Next, using a magnetic field pressing device, the alloy powder is formed while orienting the easy magnetization axis of the alloy powder in an applied magnetic field to obtain a compacted powder compact. The forming is preferably carried out in a vacuum, a nitrogen gas atmosphere, an inert gas atmosphere such as Ar, etc. to suppress oxidation of the alloy powder. The step of sintering the compacted powder compact is carried out at a temperature of 800°C or higher and 1200°C or lower in a vacuum or an inert atmosphere using a sintering furnace. If it is less than 800°C, sintering hardly progresses and a high sintering density cannot be obtained. If it exceeds 1200°C, the main phase of the Nd 2 Fe 14 B-type compound decomposes and α-Fe precipitates. The sintering temperature is particularly preferably in the range of 900 to 1100°C. The sintering time is preferably 0.5 to 20 hours, more preferably 1 to 10 hours. The sintering may be in a pattern of maintaining a constant temperature after heating, or a two-stage sintering pattern of heating to a first sintering temperature and then maintaining at a lower second sintering temperature for a predetermined time to refine the crystal grains may be used. Also, sintering may be carried out multiple times, or a discharge plasma sintering method etc. may be applied. The cooling rate after sintering is not particularly limited, but it can be cooled at a cooling rate of at least 600°C or lower, preferably 200°C or lower, preferably 1°C / min or more and 100°C / min or less, more preferably 5°C / min or more and 50°C / min or less. In order to improve the room temperature coercive force and the temperature characteristics of the coercive force, it is preferable to further carry out an aging heat treatment at 300 to 800°C for 0.5 to 50 hours. After the aging heat treatment, it can be cooled at a cooling rate of at least 200°C or lower, preferably 100°C or lower, preferably 1°C / min or more and 100°C / min or less, more preferably 5°C / min or more and 50°C / min or less. The aging heat treatment may be carried out multiple times. Also, an intermediate heat treatment at 600 to 1000°C for 0.5 to 50 hours may be carried out between the sintering heat treatment and the aging heat treatment.

[0042] The main phase grains and R(Fe, Co) 2 In order to form a thin boundary phase between the grain boundary phase, after the intermediate heat treatment, it is preferable to cool at a cooling rate of 1 °C / min or more and 50 °C / min or less, preferably 2 °C / min or more and 30 °C / min or less, to at least 550 °C or less, preferably 400 °C or less.

[0043] By performing the above intermediate heat treatment and aging heat treatment under optimal conditions according to the composition, powder particle size, etc., an R-rich phase and R(Fe, Co) 2 phase are formed at the grain boundary. In a more preferable case, a two-particle grain boundary phase is formed between adjacent main phase grains, and further, a thin boundary phase is formed between the R(Fe, Co) 2 phase and the main phase grains. This brings about an increase in the room temperature coercive force and an improvement in the temperature characteristics of the coercive force. The sintered body is cut and ground into a predetermined shape and magnetized to obtain a sintered magnet.

[0044] As shown in FIG. 1, in the sintered magnet by the two-alloy method, a main phase mainly composed of Nd 2 Fe 14 B-type compound is formed, and an R'-rich phase, R'(Fe, Co) 2 phase and the outer shell portion of the main phase grains 10 are mainly formed by the components of alloy B. Therefore, the Ce / R' atomic ratio of the R'-rich phase and R'(Fe, Co) 2 phase formed at the grain boundary portion 20 is higher than the Ce / R' atomic ratio inside the main phase grains. Also, a part of Ce in the grain boundary portion 20 substitutes for R' atoms in the surface layer portion of the main phase grains 10, forming a core-shell structure with different Ce concentrations at the center and the outer shell.

[0045] On the other hand, in the grain boundary diffusion method, first, a sintered body is produced by the single alloy method or the two alloy method in the same manner as described above. At this time, it is preferable that R' in the sintered body composition does not contain Ce.

[0046] Next, grain boundary diffusion of Ce is performed on the obtained sintered body. After cutting and grinding the sintered body as necessary, a diffusion material selected from Ce-containing compounds such as metals, alloys, oxides, fluorides, oxyfluorides, hydrides, carbides, etc. containing Ce is placed on its surface in the form of powder, thin film, ribbon, foil, etc. For example, the powder of the above material may be mixed with water or an organic solvent to form a slurry, which is coated on the sintered body and then dried, or the above substance may be placed on the surface of the sintered body as a thin film by means such as vapor deposition, sputtering, CVD, etc. The installation amount is preferably 10 to 1000 μg / mm 2 and particularly preferably 20 to 500 μg / mm 2 . In such a range, an increase in H cJ can be sufficiently obtained, and a decrease in B r due to Ce can be reduced.

[0047] This sintered body is heat-treated in a vacuum or an inert gas atmosphere with Ce placed on its surface. The heat treatment temperature is preferably 600°C or higher and the sintering temperature or lower, and particularly preferably 700°C or higher and 1000°C or lower. The heat treatment time is preferably 0.5 to 50 hours, and particularly preferably 1 to 20 hours. The cooling rate after heat treatment is not particularly limited, but 1 to 20°C / min, particularly 2 to 10°C / min is preferred. The Ce placed on the sintered body penetrates into the interior of the sintered body through the grain boundary portion by this diffusion heat treatment. At this time, as shown in FIG. 2, the R' atoms in the surface layer portion of the main phase grains 10 are replaced by Ce, and a core-shell structure with different Ce / R' ratios is formed between the central portion and the outer shell portion of the main phase grains 10, and an R'-rich phase containing Ce or R'(Fe,Co) 2 phase is formed in the grain boundary portion 20, and H cJ increases.

[0048] The sintered body subjected to the diffusion heat treatment is preferably further subjected to an aging heat treatment at 300 to 800°C for 0.5 to 50 hours in the same manner as in the case of the two-alloy method in order to improve the room temperature coercive force and the temperature characteristics of the coercive force.

[0049] The main phase grains and R(Fe,Co) 2In order to form a thin boundary phase between grain boundary phases, an intermediate heat treatment similar to that in the case of the two-alloy method may be performed on the sintered body after the diffusion treatment. In this case, it can also be omitted by combining it with the diffusion heat treatment. By performing an optimal heat treatment according to the sintered body composition, powder particle size, diffusion material, etc., an R-rich phase and R(Fe,Co) 2 phases are formed at the grain boundaries, and furthermore, a thin boundary phase is formed between the R(Fe,Co) 2 phase and the main phase grains. In a more preferred case, an intergranular phase between two particles is formed between adjacent main phase grains, and an increase in the room temperature coercive force and an improvement in the temperature characteristics of the coercive force are achieved.

[0050] Also, for further improvement of magnetic properties, Dy or Tb may be provided on the surface of this sintered body separately or simultaneously with Ce, and a diffusion heat treatment may be performed.

[0051] The anisotropic rare earth sintered magnet of the present invention thus produced has a residual magnetic flux density B of at least 12 kG or more at room temperature r and a coercive force H of 10 kOe or more cJ . Also, the temperature coefficient β of the coercive force exhibits a characteristic of β ≧ (0.01 × H cJ(室温) - 0.720)% / K. Here, β = ΔH cJ / ΔT × 100 / H cJ(室温) , (ΔH cJ = H cJ(室温) - H cJ(140℃) , ΔT = room temperature - 140 (°C)). It is more preferable that β ≧ (0.01 × H cJ(室温) - 0.7)% / K. The anisotropic rare earth sintered magnet of the present invention has a smaller temperature change in the coercive force compared to a Ce-free Nd-Fe-B sintered magnet and is suitable for use at high temperatures.

Examples

[0052] Hereinafter, examples and comparative examples will be shown to specifically explain the present invention, but the present invention is not limited to the following examples.

[0053] [Example 1] Using Nd metal, Pr metal, electrolytic iron, Co metal, ferroboron, Al metal, and Cu metal, the composition was adjusted to be 10.6 atomic % Nd, 2.7 atomic % Pr, 1.0 atomic % Co, 6.0 atomic % B, 0.5 atomic % Al, 0.1 atomic % Cu, and the balance Fe. After melting in an Ar gas atmosphere using a high-frequency induction furnace, strip casting was performed on a water-cooled Cu roll rotating at a peripheral speed of 2 m / sec, thereby producing an alloy ribbon with a thickness of about 0.2 to 0.4 mm. After polishing and etching the cross-section of this alloy, microstructure observation was carried out using a laser microscope (LEXT OLS4000 manufactured by Olympus Corporation). The observed location was set at a position about 0.15 mm from the surface where the ribbon contacted the cooling roll, and 20 observations were made. For each image, 20 lines parallel to the roll contact surface were drawn at equal intervals, and the intersections of these lines with the grain boundary phase removed by etching were counted. When calculating the average grain boundary phase interval, it was 4.7 μm. After performing a hydrogen storage treatment on this alloy at room temperature, a dehydrogenation treatment of heating at 400 °C in a vacuum was carried out to obtain coarse powder (designated as Sample 1A powder). Next, using Ce metal and electrolytic iron as raw materials, an alloy ingot with a composition adjusted to be 33 atomic % Ce and the balance Fe was manufactured using a high-frequency induction furnace. After heat treatment at 870 °C for 20 hours, it was made into coarse powder by mechanical pulverization (designated as Sample 1B powder). After mixing Sample 1A powder and Sample 1B powder at a weight ratio of 93:7, they were pulverized using a jet mill in a nitrogen gas stream to obtain fine powder with an average particle size of 3.1 μm. Next, the fine powder was filled into the mold of a molding device in an inert gas atmosphere, and while being oriented in a magnetic field of 15 kOe (= 1.19 MA / m), it was pressure-molded at a pressure of 0.6 Ton / cm 2 in the direction perpendicular to the magnetic field. The obtained compacted body was sintered in a vacuum at 1040 °C for 3 hours, cooled to room temperature, taken out once, and further heat-treated at 510 °C for 2 hours to obtain the sintered body sample of Example 1.

[0054] The obtained sintered body sample was analyzed by inductively coupled plasma optical emission spectrometry (ICP-OES) using an inductively coupled plasma optical emission spectrometer (SPS3520UV-DD manufactured by Hitachi High-Technologies Corporation). As a result, the composition was Nd 9.9 Pr 2.5 Ce 1.8 Febal. Co 1.0 B 5.6 Al 0.5 Cu 0.1 It was. From the X-ray diffraction measurement of the powder obtained by pulverizing a part of the sample, it was confirmed that the crystal structure of the main phase was Nd 2 Fe 14 type. When the structure of the sintered body was observed and the composition analysis of each phase was performed using an EPMA apparatus (manufactured by JEOL Ltd., JXA-8500F), a core / shell structure with different compositions was formed between the central part and the outer shell part of the main phase grains. Ce was not contained in R' corresponding to the core in the central part, and R' in the outer shell part of the grains contained Ce. In addition, it was confirmed that the R'-rich phase and the R'(Fe,Co) 2 phase each existed in an amount of 1% by volume or more. The volume ratio of each phase was calculated as being equal to the area ratio in the image of the backscattered electron image. The α-Fe phase and the R' 2 (Fe,Co,M) 17 phase was not observed. Since there were also phases such as oxides, the total of the phase ratios was less than 100%. Based on the analysis value of the R'(Fe,Co) 2 phase, an alloy with the same composition was prepared by arc melting, and after homogenization treatment at 800 °C for 10 hours, magnetization-temperature measurement was performed using a VSM. As a result, the Curie temperature T c was 66 °C.

[0055] As a result of observing the sintered body sample after etching, the average crystal grain size of the main phase calculated as described above was 4.3 μm. When the magnetic properties were measured with a B-H tracer, at room temperature, B r was 14.0 kG and H cJ was 13.6 kOe. Also, the temperature coefficient β of H cJ was -0.575% / K. Table 1 shows the ICP composition analysis values, average crystal grain size, and crystal structure of the main phase of the sintered body. Table 2 shows the conditions of the sintering heat treatment and the aging heat treatment, and the results of the magnetic properties measured with a B-H tracer. Table 3 shows the composition analysis values of each phase measured by EPMA.

[0056] [Comparative Example 1] The composition was adjusted using Nd metal, Pr metal, Ce metal, electrolytic iron, Co metal, ferroboron, Al metal, and Cu metal to produce strip-cast alloy ribbons. The average intergranular phase spacing calculated from the cross-sectional image of this alloy was 4.4 μm. After subjecting this alloy to a hydrogen storage treatment and a dehydrogenation treatment of heating at 400 °C in a vacuum to obtain coarse powder, it was pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.1 μm. After forming a compacted body by pressure molding in a magnetic field, sintering was performed at 1040 °C for 3 hours in a vacuum, cooled to room temperature and taken out once, and further heat-treated at 510 °C for 2 hours to obtain a sintered body sample of Comparative Example 1.

[0057] From ICP analysis, the composition of the sintered body of Comparative Example 1 was Nd 10.0 Pr 2.6 Ce 1.8 Fe bal. Co 1.0 B 5.6 Al 0.4 Cu 0.1 It was. The main phase was Nd 2 Fe 14 It was confirmed by X-ray diffraction that it had a B-type crystal structure. When microstructure observation and composition analysis of each phase were performed with an EPMA apparatus, the composition within the main phase grains was almost uniform, and there was no difference in the Ce concentration between the central part and the outer shell part. Also, although an R'-rich phase existed at the grain boundary, the R'(Fe,Co) 2 phase could not be confirmed. The average crystal grain size of the main phase was 4.0 μm. The magnetic properties measured with a B-H tracer were B r 13.7 kG and H cJ 9.8 kOe at room temperature, and the temperature coefficient β of H cJ was -0.641% / K. The results are shown in Tables 1 to 3.

[0058] [Example 2, Comparative Example 2] In Example 2, similar to Example 1, a strip-cast alloy ribbon with a composition of 12.8 atomic% Nd, 1.0 atomic% Co, 5.9 atomic% B, 0.2 atomic% Al, 0.05 atomic% Zr, and the balance being Fe, having a thickness of about 0.2 to 0.4 mm and an average grain boundary phase interval of 3.9 μm, was produced. Hydrogen storage treatment and dehydrogenation treatment were performed to obtain coarse powder (Powder of Example 2A). On the other hand, an alloy adjusted to have a composition of 80 atomic% Ce, 10 atomic% Cu, and the balance being Fe was melted in a quartz tube using a high-frequency induction furnace and sprayed onto a Cu roll rotating at a peripheral speed of 23 m / sec to produce a rapidly solidified alloy ribbon with a thickness of about 100 to 250 μm. This alloy ribbon was pulverized into coarse powder (Powder of Example 2B) by ball milling. After mixing Powder of Example 2A and Powder of Example 2B at a weight ratio of 96:4, they were pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 2.8 μm. After forming a compacted body by pressure molding in a magnetic field, sintering was performed in a vacuum at 1020 °C for 2 hours, cooled to room temperature and taken out once, and further heat-treated at 530 °C for 4 hours to obtain a sintered body sample of Example 2.

[0059] In Comparative Example 2, a strip-cast alloy ribbon with a composition of 7.8 atomic% Nd, 5.0 atomic% Ce, 1.0 atomic% Co, 5.9 atomic% B, 0.2 atomic% Al, 0.05 atomic% Zr, and the balance being Fe, having a thickness of about 0.2 to 0.4 mm and an average grain boundary phase interval of 4.2 μm, was produced. Hydrogen storage treatment and dehydrogenation treatment were performed to obtain coarse powder (Powder of Comparative Example 2A). On the other hand, an alloy adjusted to have a composition of 80 atomic% Nd, 10 atomic% Cu, and the balance being Fe was melted in a quartz tube using a high-frequency induction furnace and sprayed onto a Cu roll rotating at a peripheral speed of 22 m / sec to produce a rapidly solidified alloy ribbon with a thickness of about 100 to 250 μm. This alloy ribbon was pulverized into coarse powder (Powder of Comparative Example 2B) by ball milling. After mixing Powder of Comparative Example 2A and Powder of Comparative Example 2B at a weight ratio of 96:4, they were pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 2.8 μm. After forming a compacted body by pressure molding in a magnetic field, sintering was performed in a vacuum at 1020 °C for 2 hours, cooled to room temperature and taken out once, and further heat-treated at 530 °C for 4 hours to obtain a sintered body sample of Comparative Example 2.

[0060] The sintered body compositions of Example 2 and Comparative Example 2 were Nd respectively by ICP analysis. 12.4Ce 1.7 Fe bal. Co 1.0 B 5.7 Al 0.1 Cu 0.2 Zr 0.1 and Nd 9.2 Ce 4.9 Fe bal. Co 0.9 B 5.8 Al 0.1 Cu 0.2 Zr 0.1 It was. When the structure was observed, in Example 2, the main phase grains containing Ce were not present in the central part, but many main phase grains containing Ce were present in the outer shell part of the grains. Also, in the grain boundary part, the R'-rich phase and R'(Fe,Co) 2 phase each existed in an amount of 1% by volume or more. Based on the analysis value of the R'(Fe,Co) 2 phase, the T c of the alloy of the same composition prepared by arc melting was 74°C. On the other hand, in Comparative Example 2, Ce was contained in both the central part and the outer shell part of the main phase grains, and the Ce / R' ratio was higher in the central part of the grains than in the outer shell part of the grains. Also, in the grain boundary part, the R'(Fe,Co) 2 phase and the R'Cu 2 phase were formed, and the R'-rich phase could not be confirmed. The average crystal grain size of the main phase was 3.8 μm in Example 2 and 3.6 μm in Comparative Example 2. The results are shown in Tables 1 to 3. Example 2 was superior to Comparative Example 2 in both the room temperature magnetic properties and the temperature properties of H cJ .

[0061] [Examples 3 to 5] In Example 3, a strip cast alloy adjusted to have a composition of 13.0 atomic% Nd, 6.1 atomic% B, and the balance Fe, and an alloy adjusted to have a composition of 70 atomic% Ce, 5 atomic% La, 6 atomic% Ni, and the balance Al and arc melted were prepared, and they were mixed as coarse powders at a weight ratio of 94:6 in the same manner as in Example 1. The compacted powder compact prepared by jet mill pulverization and pressure molding in a magnetic field was sintered in a vacuum at 1010°C for 3 hours. Thereafter, aging heat treatment was performed at 480°C for 1 hour to obtain a sintered body sample.

[0062] Example 4 involved preparing a strip-cast alloy with a composition adjusted to 12.8 atomic% Nd, 6.0 atomic% B, 0.5 atomic% Al, 0.2 atomic% Cr, 0.3 atomic% Ti, and the balance being Fe, and a cast alloy with a composition adjusted to 28 atomic% Ce, 7 atomic% Gd, 30 atomic% Co, and the balance being Fe. These were mixed as coarse powders in a weight ratio of 90:10 in the same manner as in Example 1. The compacted powder compact prepared by jet milling and hot-pressing in a magnetic field was sintered in a vacuum at 1030 °C for 1.5 hours. The obtained sintered body was heat-treated at 900 °C for 1 hour, cooled to 500 °C or lower at a cooling rate of 3.8 °C / min, and then subjected to an aging heat treatment at 600 °C for 3 hours to obtain a sintered body sample.

[0063] Example 5 involved preparing a strip-cast alloy with a composition adjusted to 13.0 atomic% Nd, 6.0 atomic% B, and the balance being Fe, and an alloy prepared by arc melting with a composition adjusted to 56 atomic% Ce, 9 atomic% Y, 10 atomic% Si, 8 atomic% Ga, and the balance being Co. These were mixed as coarse powders in a weight ratio of 95:5 in the same manner as in Example 1. The compacted powder compact prepared by jet milling and hot-pressing in a magnetic field was sintered in a vacuum at 1060 °C for 2 hours. The obtained sintered body was heat-treated at 960 °C for 2 hours, cooled to 500 °C or lower at a cooling rate of 4.5 °C / min, and then subjected to an aging heat treatment at 680 °C for 3 hours to obtain a sintered body sample.

[0064] The results of Examples 3 to 5 are shown in Tables 1 to 3. In all of the sintered body microstructures, there were many primary phase grains containing Ce in the outer shell of the grains and not containing Ce in the center of the grains. At the grain boundaries, the total volume of the R'-rich phase and the R'(Fe,Co) 2 phase was present at 1% by volume or more. Also, all of the magnetic properties had an H cJ at room temperature of 10 kOe or more, and a temperature coefficient β of H cJ of (0.01 × H cJ(室温) - 0.720)% / K or more, showing good magnetic properties.

[0065] [Example 6, Comparative Example 3] The composition was adjusted using Nd metal, electrolytic iron, Co metal, ferroboron, and Al metal, and strip-cast alloy ribbons were produced. The average grain boundary spacing calculated from the cross-sectional image of this alloy was 4.8 μm. This alloy was subjected to a hydrogen storage treatment and a dehydrogenation treatment by heating at 400 °C in a vacuum to form coarse powder, which was then pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.5 μm. It was formed into a compacted powder body by pressure molding in a magnetic field and sintered at 1040 °C for 3 hours in a vacuum. Furthermore, the obtained sintered body was machined by cutting to a size of 10 × 10 × 3 mm. Next, an alloy ingot with a composition adjusted to 25 atomic% Ce, 8 atomic% Dy, 30 atomic% Co, 10 atomic% Cu, and the balance Fe was produced using Ce metal, Dy metal, electrolytic iron, Co metal, and Cu metal as raw materials in a high-frequency induction furnace. After heat treatment at 420 °C for 20 hours, it was pulverized with a ball mill to obtain powder with an average particle size of 14.6 μm. The above sintered body was immersed in a liquid obtained by mixing and stirring this powder and ethanol at a weight ratio of 1:1, pulled out, and dried by blowing to apply the powder to the surface of the sintered body. This sample was subjected to a diffusion heat treatment at 870 °C for 10 hours in a vacuum and then cooled to 500 °C or lower at a cooling rate of 5 °C / min, and further subjected to an aging heat treatment at 560 °C for 2 hours in an Ar gas atmosphere to obtain a sintered body sample of Example 6. On the other hand, a sample that was not subjected to the above powder coating and diffusion heat treatment but only subjected to an aging heat treatment at 560 °C for 2 hours in an Ar gas atmosphere was used as a sintered body sample of Comparative Example 3.

[0066] From ICP analysis, the sintered body compositions of Example 6 and Comparative Example 3 were Nd 13.6 Dy 0.1 Ce 0.6 Fe bal. Co 1.2 B 5.8 Al 0.2 Cu 0.1 , Nd 14.0 Fe bal. Co 0.4 B 6.0 Al 0.1 respectively. From EPMA microstructure observation at a depth of 500 μm from the surface of the sintered body, in Example 6, there were many primary phase grains containing Ce in the outer grain shell and no Ce in the central part, and an R'-rich phase and R'(Fe,Co) were present at the grain boundaries.2 The phases each existed in an amount of 1% by volume or more. R’(Fe, Co) 2 Based on the analysis values of the phases, the T of an alloy of the same composition prepared by arc melting c was 131°C. On the other hand, Comparative Example 3 did not contain Ce, and an R’-rich phase existed at the grain boundary, but the R’(Fe, Co) 2 phase could not be confirmed. The average crystal grain size of the main phase was 4.6 μm in both Example 6 and Comparative Example 3. The results are shown in Tables 1, 2, and 4. Example 6 showed better H cJ temperature characteristics than Comparative Example 3.

[0067] [Examples 7 to 9] In Example 7, using Nd metal, Pr metal, electrolytic iron, Co metal, ferroboron, Al metal, pure silicon, and Nb metal, the composition was adjusted to be 11.6 atomic% Nd, 2.9 atomic% Pr, 5.7 atomic% B, 1.0 atomic% Co, 0.3 atomic% Al, 0.3 atomic% Si, 0.5 atomic% Nb, and the balance Fe, and a strip-cast alloy ribbon was produced. The average grain boundary phase interval calculated from the cross-sectional image of this alloy was 4.4 μm. This alloy was subjected to a hydrogen storage treatment and a dehydrogenation treatment of heating at 400°C in a vacuum to obtain coarse powder, which was pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.1 μm. It was formed into a compacted powder compact by pressure molding in a magnetic field and sintered at 1040°C for 3 hours in a vacuum. The obtained sintered body was sized to 10 × 10 × 3 mm by cutting. Next, a metal Ce target with a diameter of 2 inches and a thickness of 3 mm was installed in a sputtering apparatus (manufactured by Canon Anelva Corporation, EB1000), and sputtering was performed for 40 minutes at an input power of 300 W and an Ar pressure of 0.5 Pa to form a Ce film on one 10 × 10 mm surface of the above sintered body. This sample was subjected to a diffusion heat treatment at 800°C for 15 hours in a vacuum and then cooled to 500°C or lower at a cooling rate of 5.3°C / min, and further subjected to an aging heat treatment at 550°C for 1 hour in an Ar gas atmosphere to obtain a sintered body sample of Example 7.

[0068] In Example 8, a strip-cast alloy was prepared with the composition adjusted to 14.1 atomic% Nd, 6.0 atomic% B, 0.5 atomic% Al, 0.1 atomic% Cu, and the balance Fe, and a strip-cast alloy ribbon was produced. The average grain boundary spacing calculated from the cross-sectional image of this alloy was 4.8 μm. This alloy was subjected to a hydrogen storage treatment and a dehydrogenation treatment of heating at 400 °C in a vacuum to obtain coarse powder, which was then pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.3 μm. It was formed into a compacted body by pressure molding in a magnetic field and sintered at 1030 °C for 2 hours in a vacuum. The obtained sintered body was sized to 10 × 10 × 3 mm by cutting. Next, the above sintered body was immersed and pulled out in a liquid obtained by mixing and stirring Ce oxide powder and pure water at a weight ratio of 3:2, and dried by blowing air to apply powder to the surface of the sintered body. This sample was subjected to a diffusion heat treatment at 880 °C for 20 hours in a vacuum and then cooled to 450 °C or lower at a cooling rate of 4.2 °C / min, and further subjected to an aging heat treatment at 510 °C for 2 hours in an Ar gas atmosphere to obtain a sintered body sample of Example 8.

[0069] In Example 9, a strip-cast alloy was prepared with the composition adjusted to 14.5 atomic% Nd, 1.0 atomic% Co, 6.2 atomic% B, 0.2 atomic% Al, 0.1 atomic% Cu, 0.05 atomic% Zr, and the balance Fe, and an alloy prepared by arc melting with the composition adjusted to 30 atomic% Ce, 35 atomic% Co, and the balance Fe were produced. They were mixed as coarse powder at a weight ratio of 95:5 in the same manner as in Example 1 and pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.7 μm. It was formed into a compacted body by pressure molding in a magnetic field and sintered at 1020 °C for 3 hours in a vacuum. The obtained sintered body was sized to 10 × 10 × 3 mm by cutting. Next, the above sintered body was immersed and pulled out in a liquid obtained by mixing and stirring Tb oxide powder and pure water at a weight ratio of 1:1, and dried by blowing air to apply powder to the surface of the sintered body. This sample was subjected to a diffusion heat treatment at 830 °C for 20 hours in a vacuum and then cooled to 500 °C or lower at a cooling rate of 5 °C / min, and further subjected to an aging heat treatment at 530 °C for 1.5 hours in an Ar gas atmosphere to obtain a sintered body sample of Example 9.

[0070] The results of Examples 7 to 9 are shown in Tables 1, 2, and 4. In any of the sintered body microstructures, the central portion did not contain Ce, and there were many primary phase grains containing Ce in the outer shell portion of the grains. In the grain boundary portion, the R'-rich phase and R'(Fe,Co) 2 phases each existed in an amount of 1% by volume or more. Also, all of the magnetic properties had an H cJ of 10 kOe or more at room temperature and an H cJ temperature coefficient β of (0.01×H cJ(室温) -0.720)% / K or more, indicating good magnetic properties.

[0071] [Example 10, Comparative Example 4] A strip-cast alloy ribbon with a composition of 13.5 atomic% Nd, 6.0 atomic% B, 0.5 atomic% Al, 0.2 atomic% Cu, and the balance Fe, having a thickness of about 0.2 to 0.4 mm and an average grain boundary phase interval of 4.1 μm, was produced and subjected to a hydrogen storage treatment and a dehydrogenation treatment to obtain coarse powder (Powder of Ex. 10A). Next, an alloy with a composition adjusted to 35 atomic% Ce, 10 atomic% Co, and the balance Fe was produced using an arc melting furnace, heat-treated at 850°C for 15 hours, and then mechanically pulverized to obtain coarse powder (Powder of Ex. 10B). After mixing the Powder of Ex. 10A and the Powder of Ex. 10B at a weight ratio of 92:8, they were pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.6 μm. After forming a compacted body by pressure molding in a magnetic field, sintering was performed in a vacuum at 1000°C for 2 hours, cooled to room temperature and taken out once, and further heat-treated at 500°C for 3 hours to obtain a sintered body sample of Example 10. On the other hand, a sample produced in the same manner as in Example 10 up to the sintering step was heat-treated at 980°C for 1 hour and then cooled in an Ar atmosphere, which was used as Comparative Example 4.

[0072] From ICP analysis, the sintered body compositions of Example 10 and Comparative Example 4 were Nd 12.5 Ce 2.1 Fe bal. Co 0.7 B 5.8 Al 0.4 Cu 0.1 . In the EPMA microstructure observation, neither contained Ce in the central portion, and there were many primary phase grains containing Ce in the outer shell portion of the grains. In Example 10, the R'-rich phase and R'(Fe,Co) were present in the grain boundary portion2 The phases each existed in an amount of 1% by volume or more. R’(Fe, Co) 2 Based on the analytical values of the phases, the T of an alloy of the same composition produced by arc melting c was 70°C. On the other hand, in Comparative Example 4, an R’-rich phase was present at the grain boundaries, but the R’(Fe, Co) 2 phase could not be confirmed. The average crystal grain size of the main phase was 4.9 μm in both Example 10 and Comparative Example 4. The results are shown in Tables 1, 2, and 4. Example 10 had a higher room-temperature H cJ than Comparative Example 4, and the temperature characteristics of H cJ were also good.

[0073] [Example 11] A strip-cast alloy ribbon with a composition of 13.5 atomic% Nd, 5.9 atomic% B, 1.0 atomic% Co, 0.5 atomic% Al, 0.2 atomic% Cu, 0.1 atomic% Zr, and the balance Fe, having a thickness of about 0.2 to 0.4 mm and an average grain boundary phase interval of 4.2 μm, was produced and subjected to a hydrogen storage treatment and a dehydrogenation treatment to obtain coarse powder (Powder of Example 11A). Next, an alloy ingot with a composition adjusted to 33.3 atomic% Ce, 1.0 atomic% Co, and the balance Fe was produced using an arc melting furnace, heat-treated at 860°C for 18 hours, and then mechanically pulverized to obtain coarse powder (Powder of Example 11B). After mixing Powder of Example 11A and Powder of Example 11B at a weight ratio of 93:7, they were pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 2.9 μm. After forming a compacted body by pressure molding in a magnetic field, sintering was performed in a vacuum at 1020°C for 3 hours, cooled to room temperature, and taken out once. Next, an intermediate heat treatment was performed in an Ar atmosphere at 900°C for 1 hour, then cooled to 450°C or lower at a cooling rate of 5°C / min, and subsequently, a low-temperature heat treatment was performed at 510°C for 3 hours to obtain a sintered body sample of Example 11.

[0074] The composition of this sintered body was Nd 12.7 Ce 1.8 Fe bal. Co 1.1 B 5.6 Al 0.5 Cu 0.1 Zr 0.1It was the case. In the EPMA structure observation, there were many primary phase grains that did not contain Ce in the central part and contained Ce in the outer shell part of the grains. Also, an R'-rich phase and an R'(Fe,Co) 2 phase each existed in an amount of 1% by volume or more. Based on the analysis values of the R'(Fe,Co) 2 phase, the T of an alloy of the same composition prepared by arc melting was c 68 °C. The average crystal grain size of the primary phase was 3.9 μm. The results are shown in Tables 1, 2, and 5.

[0075] Using an FIB-SEM apparatus (Scios manufactured by FEI Company), an observation sample was cut out from the sample of Example 11 and observed with a STEM apparatus (JEM-ARM200F manufactured by JEOL Ltd.). As shown in the HAADF image of Fig. 5, it was confirmed that a boundary phase was formed between the R'(Fe,Co) 2 phase at the grain boundary and the primary phase. The average thickness of this boundary phase was 1.4 nm, and the composition of the boundary phase measured by EDS analysis was Nd 22.5 Ce 13.5 Fe bal. Co 3.0 Cu 1.7 . On the other hand, the EDS analysis composition of the adjacent R'(Fe,Co) 2 phase was Nd 14.7 Ce 19.5 Fe bal. Co 2.3 Cu 0.1 . From this, it can be seen that this boundary phase is a phase having a composition different from that of the R'(Fe,Co) 2 phase. At another location of the same sample, an inter-particle grain boundary phase with an average thickness of about 2.4 nm existed between adjacent primary phase grains, and its average composition was Nd in terms of the EDS analysis value 26.8 Ce 6.9 Fe bal. Co 7.4 Cu 12.5 Zr 0.5 . From this, when calculating Ce / R' for the boundary phase formed between the primary phase and the R'(Fe,Co) 2 phase and the inter-particle grain boundary phase between the primary phase grains, they are 0.37 and 0.20 respectively, and it can be seen that the former shows a higher Ce / R'.

[0076] [Example 12] A strip-cast alloy ribbon with a composition of 10.6 atomic% Nd, 2.5 atomic% Pr, 5.9 atomic% B, and the balance Fe, a thickness of about 0.2 to 0.4 mm, and an average grain boundary spacing of 4.0 μm was fabricated. After hydrogen storage treatment and dehydrogenation treatment, it was pulverized with a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 3.0 μm. It was made into a compacted body by pressure molding in a magnetic field, and the sintered body obtained by sintering in a vacuum at 1040 °C for 2 hours was sized to 10 × 10 × 3 mm by cutting. Next, a target with a composition of Ce 30 Fe bal. Co 20 Al 20 Cu 5 V 5 with a diameter of 2 inches and a thickness of 3 mm was used, and sputtering was performed for 90 minutes with an input power of 250 W and an Ar pressure of 0.4 Pa to form a Ce film on one 10 × 10 mm surface of the above sintered body. After this sample was subjected to diffusion heat treatment in a vacuum at 840 °C for 25 hours, it was cooled to 500 °C or lower at a cooling rate of 4.5 °C / min, and then further subjected to aging heat treatment at 540 °C for 3 hours in an Ar gas atmosphere to obtain a sintered body sample of Example 12.

[0077] The sintered body composition of Example 12 was Nd 10.2 Pr 2.4 Ce 1.0 Fe bal. Co 0.6 B 5.6 Al 0.2 Cu 0.1 V 0.1 as determined by ICP analysis. In the EPMA tissue observation, there were many main phase grains containing Ce in the outer grain shell but not in the central part. Also, in the grain boundary part, the R'-rich phase and R'(Fe,Co) 2 phase each existed in an amount of 1 volume% or more. The T 2 of an alloy with the same composition prepared by arc melting based on the analysis value of the R'(Fe,Co) c phase was 78 °C. The results are shown in Tables 1, 2, and 5.

[0078] STEM observation was performed on the structure of Example 12, and R'(Fe,Co)2 A boundary phase with an average thickness of 1.6 nm and a composition of Nd 20.1 Pr 2.6 Ce 13.7 Fe bal. Co 2.5 Cu 1.9 was confirmed to be formed between the phase and the main phase. From this, the Ce / R’ of the boundary phase was calculated to be 0.38. On the other hand, at another location of the same sample, there was an intergranular phase with an average thickness of about 1.8 nm between adjacent main phase grains, and its average composition was Nd 17.7 Pr 6.2 Ce 6.9 Fe bal. Co 7.3 Cu 8.9 V 0.4 . (Ce / R’ = 0.22) From this, it can be seen that the Ce / R’ of the boundary phase formed between the main phase and the R’(Fe,Co) 2 phase is higher than that of the intergranular phase between two particles.

[0079]

Table 1

[0080]

Table 2

[0081]

Table 3

[0082]

Table 4

[0083]

Table 5

Explanation of Symbols

[0084] 11 Main phase (region with high Ce / R’) 12 Matrix phase (region with low Ce / R') 21 R'-rich phase 22 R'(Fe, Co) 2 phase 31 Inter-particle grain boundary phase formed between adjacent matrix phase grains 32 R'(Fe, Co) 2 Boundary phase formed between the R'(Fe, Co) phase and the matrix phase

Claims

1. The composition is represented by the formula R ｘ (Fe １－ａ Co ａ ) １００－ｘ－ｙ－ｚ B ｙ M ｚ (R is at least two elements selected from rare earth elements and including Nd and Ce as essential elements, M is at least one element selected from the group consisting of Al, Si, Ti, V, Cr, Mn, Ni, Cu, Zn, Ga, Ge, Zr, Nb, Mo, Ag, In, Sn, Hf, Ta, W, Pb, Bi, and x, y, z, a are each 12 ≦ x ≦ 17 atomic %, 3.5 ≦ y ≦ 6.0 atomic %, 0 ≦ z ≦ 3 atomic %, 0 ≦ a ≦ 0.1). The anisotropic rare earth sintered magnet is characterized in that the main phase is composed of a Nd ２ Fe １４ B-type crystal compound, and there are main phase grains in which the Ce / R' ratio (R' is at least one element selected from rare earth elements and including Nd as essential element) in the central part of the grains is lower than the Ce / R' ratio in the outer shell part of the grains, and there are an R'-rich phase containing Ce and an R'(Fe, Co) ２ phase containing Ce at the grain boundary.

2. between the main phase and the R'(Fe, Co) ２ The anisotropic rare earth sintered magnet according to claim 1, wherein a boundary phase containing 20 atomic% or more of R and having a thickness of 20 nm or less is formed between the main phase and the R'(Fe, Co) phase.

3. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein in the main phase grains, there are main phase grains in which Ce is not contained in R' at the center part.

4. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein in the main phase grains, there are main phase grains in which R' at the center part is Nd, or consists of Nd and Pr.

5. The above R'(Fe, Co) ２ The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein the phase is a phase exhibiting ferromagnetic or ferrimagnetic properties at room temperature or higher.

6. said R'(Fe, Co) ２ The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the Ce / R' ratio in the phase is higher than the Ce / R' ratio in the outer shell portion of the main phase grains.

7. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein the Ce / R' ratio in the R'-rich phase is higher than the Ce / R' ratio in the outer shell part of the main phase grains.

8. The R'-rich phase and R'(Fe, Co) ２ The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the phases are contained in a total amount of 1% by volume or more.

9. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein the Ce / R' ratio in the composition of the sintered body is 0.01 or more and 0.3 or less.

10. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein the B-rich phase contained in the sintered magnet is 5% by volume or less.

11. The anisotropic rare earth sintered magnet according to claim 1 or 2, wherein a two-particle grain boundary phase is formed between adjacent main phase grains.

12. The Ce / R' in the boundary phase formed between the main phase and the R'(Fe, Co) ２ phase is higher than the Ce / R' in the two-particle grain boundary phase formed between adjacent main phase grains, and the anisotropic rare earth sintered magnet according to claim 11, characterized in that.

13. Coercivity H at room temperature ｃＪ（ｒｏｏｍ） is 10 kOe or more, and the value of the temperature coefficient β of the coercivity satisfies β ≧ (0.01 × H ｃＪ（室温） − 0.720)% / K. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized by this.

14. Nd ２ Fe １４ A method for producing an anisotropic rare-earth sintered magnet according to claim 1 or 2, characterized in that an alloy containing a compound phase of a B-type crystal and an alloy having a higher R' composition ratio and Ce / R' ratio than that are pulverized and mixed, and then compacted under a magnetic field application to form a compact, and then sintered at a temperature of 800°C or higher and 1200°C or lower.

15. Nd ２ Fe １４ An alloy containing a compound phase of B-type crystal is pulverized, compacted under the application of a magnetic field to form a compact, and then sintered at a temperature of 800°C or higher and 1200°C or lower. A material containing Ce is brought into contact with the sintered body, and heat treatment is performed at a temperature of 600°C or higher and lower than the sintering temperature to diffuse Ce into the interior of the sintered body. The method for producing an anisotropic rare-earth sintered magnet according to claim 1 or 2, characterized by the above steps.

16. The method for manufacturing an anisotropic rare earth sintered magnet according to claim 15, wherein the material containing Ce brought into contact with the sintered body is one or more selected from Ce metal, Ce-containing alloy, and Ce-containing compound, and its form is one or more selected from powder, thin film, thin strip, foil, and gas.

17. The method for manufacturing an anisotropic rare earth sintered magnet according to claim 14, wherein the sintered body is heat-treated at a temperature of 300 to 800°C.

18. The method for manufacturing an anisotropic rare earth sintered magnet according to claim 14, wherein after heat-treating the sintered body at a temperature of 600 to 1000°C, it is cooled at a cooling rate of 1°C / min or more and 50°C / min or less to at least 550°C or lower, and then heat-treated at a temperature of 300 to 800°C.

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