JP7684255B2 - Anisotropic rare earth sintered magnet and its manufacturing method - Google Patents

Description

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

Demand for Nd-Fe-B sintered magnets is expected to grow further in the future, driven by the electrification of automobiles and the trend toward higher performance and energy saving in industrial motors. However, the rare earth elements used as raw materials, such as Nd, Pr, Dy and Tb, are expensive and pose a risk to future supply stability. For this reason, research is being conducted to replace some of the Nd with Ce, which is an inexpensive element with a higher concentration in the earth's crust.

For example, in Patent Document 1, a steel sheet is disclosed that includes a main phase and a grain boundary phase, and has an overall composition represented by (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., and M ² is a transition metal element that is alloyed with R ³ , etc.), and the main phase is R ₂ Fe ₁₄ The present invention discloses a rare earth magnet having excellent coercivity and residual magnetization, the magnet having a B-type crystal structure, an average grain size of the main phase of 1 to 20 μm, the main phase having a core and a shell, the thickness of the shell being 25 to 150 nm, and satisfying 0≦b≦0.30 and 0≦b/a≦0.50, where a is the light rare earth element ratio in the core and b is the light rare earth element ratio in the shell. The magnet also has a manufacturing method therefor.

Patent Document 2 also discloses a rare earth magnet that has main phase particles containing R, T and B, and a grain boundary phase, where R contains Nd and Ce, T contains Fe, the grain boundary phase includes an R-T phase and an R-rich phase, the R-T phase contains intermetallic compounds of R and T, the R content in the R-rich phase is greater than the R content in the R-T phase, Ce/R x 100 = 65-100 in the R-T phase, and the R content in the R-rich phase is 70-100 atomic %.

In Patent Document 3, the overall composition is represented by _the formula (Nd _{(1-xy) CexR1y} ) _p (Fe _(1-z) ^Coz _{) (100-pq-rs) BqGarMs ( where} R The rare earth magnet has a magnetic phase and a grain boundary phase surrounding the magnetic phase, and a method for manufacturing the same is disclosed. In the rare earth magnet, p< ^0.1 is one or more elements selected from rare earth elements other than Nd and Ce and Y, M is one or more elements selected from Al, Cu, Au, Ag, Zn, In, Mn, Zr, and Ti, and unavoidable impurity elements, and is represented by the formula: 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≦z0.140.

Patent Document 4 discloses a permanent magnet having high flexural strength, which comprises a plurality of main phase particles containing a rare earth element R, a transition metal element T, and boron B, and a grain boundary phase located between the plurality of main phase particles, where R contains Nd and Ce, T contains 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 0.1-0.6, [T]/[B] is 14-18, the grain boundary phase includes an R-T phase containing an intermetallic compound of R and T, the area of a 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 0.05-0.5.

Patent Document 5 discloses a rare earth magnet comprising crystal grains having an overall composition of (Ce _x Nd _(1-x) ) _y Fe _{(100-y-w-z-v)} Co _w B _z M _v (wherein M is at least one of Ga, Al, Cu, Au, Ag, Zn, In, and Mn, and 0≦x≦0.75, 5≦y≦20, 4≦z≦6.5, 0≦w≦8, 0≦v≦2), which is composed of a core portion 1 and a surrounding shell portion 2, with the shell portion 2 having a higher Nd concentration than the core portion 1.

Patent Document 6 discloses an R-T-B magnet in which R1 and Ce are essential as R, and the raw material R-T-B magnet is subjected to a long-term heat treatment to turn the main phase particles into a core-shell. When the mass concentrations of R1 and Ce in the core portion are αNd and αCe, respectively, and the mass concentrations of R1 and Ce in the shell portion are βR1 and βCe, respectively, the ratio (B/A) of the mass concentration ratio of R1 and Ce in the shell portion (βR1/βCe=B) to the mass concentration ratio of R1 and Ce in the core portion (αR1/αCe=A) is set to 1.1 or more, thereby improving the adhesion strength with the plating by adding Ce, while suppressing a decrease in coercivity.

JP 2021-44361 A JP 2020-95989 A JP 2019-179796 A JP 2018-174323 A JP 2016-111136 A JP 2014-216339 A

As mentioned above, it has been proposed that good properties can be obtained in R-T-B magnets containing Ce by providing main phase grains with a core-shell structure or by using an R-T intermetallic compound as the grain boundary phase. However, the magnetic properties of the main phase R ₂ Fe ₁₄ B compound at room temperature are low, with a saturation magnetization M _s of 1.60 T and an anisotropic magnetic field μ ₀ H _A of 6.7 T when R=Nd, but M _s of 1.17 T and μ ₀ H _A of 3.0 T when R=Ce, making it difficult to solve the problem of reduced magnetic properties when the Ce content is high.

The present invention has been made in consideration of the above problems, and aims to provide an anisotropic rare earth sintered magnet that has _{a Nd2Fe14B} type crystal compound as a main phase and contains Ce, and that exhibits good magnetic properties, and a method for producing the same.

As a result of extensive investigations to achieve the above-mentioned object, the inventors have discovered that in an anisotropic rare earth sintered magnet containing Ce and having a main phase of an Nd ₂ Fe ₁₄ B type crystal compound, there are main phase grains in which the Ce/R' ratio (R' is one or more elements selected from rare earth elements, with Nd being an essential element) in the center of the grain is lower than the Ce/R' ratio in the outer shell of the grain, and an R'-rich phase containing Ce and an R'(Fe, Co) ₂ phase containing Ce are present in the grain boundary portions, and thus good magnetic properties can be obtained, thereby completing the present invention.

Therefore, the present invention provides the following anisotropic rare earth sintered magnet and method for producing 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 (R is two or more elements selected from rare earth elements, with Nd and Ce being essential, 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, and 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), and the main phase is Nd ₂ Fe ₁₄ An anisotropic rare earth sintered magnet is characterized in that it is made of a B-type crystal compound, has main phase grains in which the Ce/R' ratio (R' is one or more elements selected from rare earth elements, and Nd is an essential element) at the center of the grain is lower than the Ce/R' ratio at the outer periphery of the grain, and has an R'-rich phase containing Ce and an R'(Fe, Co) ₂ phase containing Ce at grain boundaries.
(2) The anisotropic rare earth sintered magnet according to ₍₁ ), characterized in that 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) two phases.
(3) The anisotropic rare earth sintered magnet according to (1) or (2), wherein the main phase grains include main phase grains in which R' at the center does not contain Ce.
(4) The anisotropic rare earth sintered magnet according to any one of (1) to (3), wherein the main phase grains include main phase grains in which R' at the center is Nd or Nd and Pr.
(5) The anisotropic rare earth sintered magnet according to any one of (1) to (4), wherein the R'(Fe, Co) ₂ phase is a phase that exhibits ferromagnetism or ferrimagnetism at room temperature or higher.
(6) The anisotropic rare earth sintered magnet according to any one of (1) to (5), wherein the Ce/R' ratio in the R'(Fe, Co) ₂ phase is higher than the Ce/R' ratio in the outer shell portion of the main phase grains.
(7) An 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 portion of the main phase grains.
(8) The anisotropic rare earth sintered magnet according to any one of (1) to (7), characterized in that 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) An anisotropic rare earth sintered magnet according to any one of (1) to (10), characterized in that a grain boundary phase is formed between adjacent main phase grains.
(12) The anisotropic rare earth sintered magnet according to ( ₁₁ ), wherein the Ce/R' in the boundary phase formed between the main phase and the R'(Fe, Co) two phases is higher than the Ce/R' in the grain boundary phase between two grains formed between adjacent main phase grains.
(13) An anisotropic rare earth sintered magnet as described in any one of (1) to (12), characterized in that the coercive force HcJ _(room) at room temperature is 10 kOe or more, and the value of the temperature coefficient of coercive force β is expressed as β≧(0.01×HcJ _{(room temperature)} −0.720)%/K.
(14) A method for producing an anisotropic rare earth sintered magnet according to any one of (1) to ( ₁₃ ), characterized in that an alloy containing a _Nd2Fe14B type crystal compound phase and an alloy having a higher R' composition ratio and Ce/R' ratio than that of the alloy are crushed and mixed, compacted into a green body by applying a magnetic field, and then sintered at a temperature of 800°C to 1200°C.
(15) A method for producing an anisotropic rare earth sintered magnet according to any one of (1) to (14), characterized in that an alloy containing a compound phase _{of Nd2Fe14B} type crystal is pulverized, compacted in an applied magnetic field to form a green body, sintered at a temperature of 800°C to 1200°C, and the sintered body is brought into contact with a material containing Ce, and heat-treated at a temperature of 600°C to the sintering temperature to diffuse Ce into the sintered body.
(16) The method for producing an anisotropic rare earth sintered magnet according to (15), characterized in that the Ce-containing material brought into contact with the sintered body is one or more selected from the group consisting of Ce metal, a Ce-containing alloy, and a Ce-containing compound, and is in the form of one or more selected from the group consisting of a powder, a thin film, a thin ribbon, a foil, and a gas.
(17) A method for producing an anisotropic rare earth sintered magnet according to any one of (14) to (16), characterized in that the sintered body is subjected to a heat treatment at a temperature of 300 to 800° C.
(18) A method for producing an anisotropic rare earth sintered magnet according to any one of (14) to (17), characterized in that the sintered body is subjected to a heat treatment at a temperature of 600 to 1000°C, then cooled to at least 550°C at a cooling rate of 1°C/min to 50°C/min, and further subjected to a heat treatment at a temperature of 300 to 800°C.

According to the present invention, it is possible to obtain an anisotropic rare earth sintered magnet that has a Nd ₂ Fe ₁₄ B type crystal compound as a main phase and contains Ce, and that exhibits excellent magnetic properties.

FIG. 1 is a schematic diagram of the structure of an example of an anisotropic sintered magnet of the present invention produced by the two-alloy method, in which an R'-rich phase and an R'(Fe, Co) ₂ phase exist at the grain boundaries. FIG. 2 is a schematic diagram of the structure of an example of an anisotropic sintered magnet of the present invention produced by the grain boundary diffusion method, in which an R'-rich phase and an R'(Fe, Co) ₂ phase exist at the grain boundaries. FIG. 3 is a schematic diagram of the structure of an example of an anisotropic sintered magnet of the present invention, in which an R'-rich phase and an R'(Fe, Co) ₂ phase are present at the grain boundaries, and _a boundary phase is formed between the main phase and the R(Fe, Co)2 phase. FIG. 4 is a HAADF image showing the boundary phase formed between the main phase and the R′(Fe,Co) ₂ phase in Example 11.

Hereinafter, an embodiment 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
The anisotropic rare earth sintered magnet is represented by the formula: Nd ₂ Fe ₁₄ B type crystal compound is the main phase, the main phase particles have particles with different Ce/R' ratios at the center and outer shell of the grain, and the grain boundary has an R'-rich phase containing Ce and an R'(Fe,Co) ₂ phase containing Ce. First, each component will be described below. R is selected from rare earth elements, and two or more elements with 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. Furthermore, R' is at least one element selected from rare earth elements, with Nd being essential.
The R'-rich phase is a phase containing more than 40 atomic % of R. The R'(Fe,Co) ₂ phase has a _MgCu2 structure and is a compound phase known as a Laves phase.

As described above, R is selected from rare earth elements and is two or more elements, with Nd and Ce being essential. 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 of Nd ₂ Fe ₁₄ B type crystal structure, which is the main phase. The content of R is 12 atomic % or more and 17 atomic % or less. It is more preferable if it is 12.5 atomic % or more and 16 atomic % or less. If it is less than 12 atomic %, the α-Fe phase precipitates and sintering is difficult, while 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. _{Since Nd2Fe14B} type compounds exhibit particularly high magnetic properties when R is Nd, the anisotropic rare earth sintered magnet of the present invention requires Nd. In addition, 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 keeping the Ce ratio in this range, an anisotropic sintered magnet can be obtained that has a high residual magnetic flux density B _r , a high coercive force H _cJ , and good H _cJ temperature characteristics.

B is also an essential element for forming a Nd ₂ Fe ₁₄ B type compound. The B content is 3.5 atomic % or more and 6.0 atomic % or less. It is more preferable that it is 5.0 atomic % or more and 5.8 atomic % or less. If it is less than 3.5 atomic %, phases that adversely affect the magnetic properties, such as the R ₂ Fe ₁₇ phase and the α-Fe phase, are precipitated, while 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.

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 have the effect of increasing _HcJ by dissolving in _{the Nd2Fe14B} type compound main phase or forming a grain boundary phase, but if contained in excess, it reduces the Br of the magnet. Therefore, when M is contained, its content is set to 3 atomic % or less in total. It is more preferable if it is 2 atomic % or less, and particularly preferable if it is 1 atomic % or less.

The anisotropic rare earth sintered magnet of the present invention has Fe as an essential element together with R and B. Furthermore, a part of Fe may be substituted with Co. Substitution with Co has the effect of increasing the Curie temperature _Tc of _{the Nd2Fe14B} type compound which is the main phase. The substitution rate of Co is 10% or less in atomic ratio. If the substitution rate exceeds 10%, _Ms decreases conversely. The ratio of Fe and Co is the remainder of R, B and M. In addition, inevitable impurities that are taken in from raw materials or mixed in during the manufacturing process, specifically H, C, N, O, F, P, S, Mg, Cl, Ca, etc. may be contained, but from the viewpoint of obtaining good magnetic properties, the content is preferably 3% by weight or less in total, more preferably 1% by weight or less. In particular, C, N, and O are preferably 1% by weight or less in total, more preferably 0.5% by weight or less, and particularly preferably 0.3% by weight or less.

Next, the phases that make up 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 made of a compound having a _Nd2Fe14B 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 in such a range, it is possible to suppress a decrease in the residual magnetic flux density B _r and a decrease in the coercive force H _cJ due to a decrease in the degree of orientation of the crystal grains. From the viewpoint of obtaining good B _r and H _cJ , the volume fraction of the main phase is preferably 80 volume % or more and less than 99 volume % of the entire magnet, and more preferably 90 volume % or more and 99 volume % or less.
The crystal grain size of the main phase can be calculated by polishing a cross section of the sintered magnet until it becomes a mirror surface, immersing it in an etching solution (such as a mixture of nitric acid + hydrochloric acid + glycerin) to selectively remove the grain boundary phase, and then observing at least 10 random points on the cross section with a laser microscope. The cross-sectional area of each particle is calculated from the obtained observation images by image analysis, and the average diameter when these are regarded as circles is taken as the average crystal grain size.
In addition, the volume fractions of the main phase and each phase can be calculated by polishing the cross section of the sintered magnet until it becomes a mirror surface, and then observing the structure of the anisotropic rare earth sintered magnet using an EPMA and analyzing the composition of each phase, and confirming the presence of the main phase, the R'-rich phase, and the R'(Fe, Co) ₂ phase, and then assuming that the area ratio in the backscattered electron image is equal to the volume fraction of each phase.

The R' ₂ Fe ₁₄ B compound has the highest saturation magnetization _Ms when R'=Nd, and when part of Nd is substituted with Ce, the larger the Ce substitution amount, the lower the _Ms. Therefore, in the magnet of the present invention, in order to reduce the effect of the decrease in B _r of the magnet due to Ce substitution, the Ce/R' ratio (atomic ratio of Ce to R') is different between the center and the outer shell of the main phase grain, and there are main phase grains in which the Ce/R' ratio in the center of the grain is lower than the Ce/R' ratio in the outer shell of the grain. However, some main phase grains may be included in which the Ce concentration distribution is uniform. Here, the outer shell refers to the region including the surface of the main phase grain, and the center refers to the other internal region. By adopting such a structure, the decrease in _Ms is suppressed in the region near the center of the main phase grain with a low Ce/R' ratio, and the decrease in B _r of the magnet due to Ce substitution can be reduced. It is more preferable that R' in the center of the main phase grain does not contain Ce, and it is even more preferable that R' in the center of the grain consists of Nd, or Nd and Pr.
On the other hand, as described later, when an R'-rich phase containing Ce and an R'(Fe, Co) ₂ phase are formed in the grain boundary, _HcJ at room temperature increases and the temperature change of _HcJ becomes small, resulting in excellent magnetic properties. In order to efficiently form these phases, the magnet of the present invention has a structure in which the Ce/R' ratio of the outer shell of the main phase grain is higher than the Ce/R' ratio of the center of the main phase grain. This increases the Ce concentration in the grain boundary, making it easier for the R'(Fe, Co) ₂ phase to form in the grain boundary. In contrast, when the Ce/R' ratio of the grains is uniform, it is necessary to increase the Ce substitution amount of the sintered body in order to significantly form the R'(Fe, Co) ₂ phase, which leads to a significant decrease in _Ms.
When the Ce/R' ratio in the outer shell of the grain is high, the H _A of the grain surface decreases, but the negative effect of the decrease in H _A is reduced because the effect of increasing H _cJ due to the Ce-containing R'-rich phase and the R'(Fe,Co) ₂ phase is large.
Conversely, if there are main phase grains in which the Ce/R' ratio in the center is higher than the Ce/R' ratio in the outer periphery of the grain, the drop in _Ms will be significant in the region near the center of the main phase grain with the high Ce/R' ratio, which is incompatible with the guidelines for the magnet of the present invention. For this reason, it is assumed that there are no main phase grains in the magnet of the present invention in which the Ce/R' ratio in the center is higher than the Ce/R' ratio in the outer periphery of the grain.
The thickness of the outer shell part having a high Ce/R' ratio is not particularly limited, but from the viewpoint of increasing the volume ratio of the inner part of the outer shell part, it is preferably 1 nm to 2 μm, and particularly preferably 2 nm to 1 μm.

The R'-rich phase and the R'(Fe,Co) ₂ phase are formed in the grain boundary of the magnet structure. The grain boundary includes the grain boundary triple point as well as the grain boundary phase between two grains. Here, 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 are present in the grain boundary, the H _cJ of the magnet at room temperature is improved, and the temperature characteristics of the 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, the R'-rich phase is difficult to exist. If it is 0.03 or more and 0.25 or less, it is more preferable, and if it is 0.05 or more and 0.2 or less, it is particularly preferable.

The R'-rich phase and the R'(Fe,Co) ₂ phase mainly bring 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 melt and become liquid phase, so liquid phase sintering proceeds and sintering is completed more quickly than solid phase sintering in which these phases are not included. In addition, by coexisting the R'-rich phase and the R'(Fe,Co) ₂ phase, the liquid phase generation temperature tends to decrease more than in the case of only one of the phases, and liquid phase sintering proceeds more quickly.

The second effect is cleaning of the main phase grain surface. Since the anisotropic rare earth sintered magnet of the present invention has a nucleation-type coercive force mechanism, it is desirable that the main phase grain surface is smooth so that nucleation of reverse magnetic domains is unlikely 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 in the sintering process or the subsequent aging process, and this cleaning effect suppresses nucleation of reverse magnetic domains, which is a factor in reducing coercive force. The R'(Fe,Co) ₂ phase has a relatively high wettability with respect to the main phase compared to other phases in which R' is less than 40 atomic %, such as compound phases such as R'M ₃ , R'M ₂ , R'(Fe,Co)M and R'(Fe,Co) ₂ M ₂ . In particular, this phase is easily able to cover the surface of the main phase grains by coexisting with the R'-rich phase, resulting in a large cleaning effect. This is believed to suppress nucleation of reverse magnetic domains, increase the coercive force at room temperature, reduce the decrease in coercive force at high temperatures, and provide good temperature dependence of _HcJ .

The third effect is the effect of weakening the magnetic interaction between main phase grains. In magnets that have _two phases, an R'-rich phase and an R'(Fe, Co) phase, an optimal sintering process or aging process is performed to form an interparticle grain boundary phase between adjacent main phase grains that contains more R' than the main phase. This weakens the magnetic interaction between the main phase grains and exerts coercive force, but if the interparticle grain boundary phase contains Ce, the effect of weakening the magnetic interaction between the main phase grains becomes even greater, and it is believed to act in the direction of further increasing the coercive force.

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 magnets in which an R'-rich phase and an R'(Fe, Co) ₂ phase exist at the grain boundary, a thin boundary phase is formed not only between the main phase grains but also between the R'(Fe, Co) ₂ phase and the main phase grains by optimally performing sintering and subsequent heat treatment according to conditions such as the composition and powder particle size. 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, resulting in a high coercive force.
In magnets that do not have an R'-rich phase at the grain boundaries, 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 particles of the main phase grains, or it is difficult for the surfaces of the main phase grains to become completely covered by these phases, making it difficult to obtain a magnet that exhibits sufficient coercivity.

As described above, the R'-rich phase contains at least 40 atomic % of R'. If the R' content exceeds 40 atomic %, the wettability with the main phase is further improved, and the above-mentioned effect is more easily obtained. It is more preferable to contain 50 atoms or more of R', and particularly preferable to contain 60 atoms or more. The R'-rich phase may be an R' metal phase, an amorphous phase, or an intermetallic compound with a high R' composition and a low melting point, such as R _{' 3} (Fe, Co, M), R' ₂ (Fe, Co, M), R' 5 (Fe, Co, M) ₃ , or R' (Fe, Co, M). It may also contain Fe, Co, M elements, and impurity elements such as H, B, C, N, O, F, P, S, Mg, Cl, and Ca, up to a total of less than 60 atomic %.
Furthermore, the higher the Ce/R ratio in the R-rich phase, the greater the effect of reducing the magnetic interaction between the main phase grains. Therefore, in order to allow Ce to efficiently act to improve the magnetic properties, it is preferable that the Ce/R ratio in the R-rich phase is higher than the Ce/R ratio in the outer shell portion of the main phase grains.

On the other hand, the R'(Fe,Co) ₂ phase is a Laves compound of _MgCu2 type crystal, but when composition analysis is performed using EPMA or the like, taking into consideration measurement variations, etc., it is assumed that it contains 20 atomic % or more and less than 40 atomic % of R'. In addition, a part of Fe and Co may be substituted with the M element. However, the amount of substitution of M is within a range in which the _MgCu2 type crystal structure is maintained.

The R'(Fe,Co) ₂ phase in the anisotropic rare earth sintered magnet of the present invention is a magnetic phase. The magnetic phase here refers to a phase that exhibits ferromagnetism or ferrimagnetism and has a Curie temperature _Tc of room temperature (23°C) or higher. _R'Fe2 has a _Tc of room temperature or higher except for _CeFe2 , and _CeFe2 also has a _Tc of room temperature or higher if 10% or more of R' is replaced with another element. On the other hand, _R'Co2 has a _Tc of room temperature or lower except for _GdCo2 , or is a paramagnetic phase, but in the anisotropic rare earth sintered magnet of the present invention, the atomic ratio of Fe substituted by Co is 0.1 or less, so in most cases the R'(Fe,Co) ₂ phase is a magnetic phase. Generally, soft magnetic phases contained in a structure often have an adverse effect on magnetic properties, but in the anisotropic rare earth sintered magnet of the present invention, the cleaning effect of the main phase grain surfaces by the R'(Fe, Co) ₂ phase and the effect of forming grain boundary phases between two particles are greater, and it is believed that even the magnetic phase contributes to an increase in room temperature _HcJ and an improvement in the temperature dependence of _HcJ .
In addition, the R'(Fe,Co) ₂ phase is unlikely to exist stably when R' is composed only of Nd and Pr, and is formed as an equilibrium phase at the grain boundary by including Ce. Therefore, it is preferable that the Ce/R' ratio of the R'(Fe,Co) ₂ phase is higher than the Ce/R' ratio of the outer shell portion of the main phase grain.

The total amount of the R'-rich phase and the R'(Fe,Co) ₂ phase is preferably 1% by volume or more, more preferably 1% by volume or more and less than 20% by volume. Moreover, it is more preferable that it is 1.5% by volume or more and less than 15% by volume, and even more preferable that it is 2% by volume or more and less than 10% by volume. Moreover, it is preferable that the R'-rich phase and the R'(Fe,Co) ₂ phase are each 0.5% by volume or more. By setting it in such a range, the area in contact with the main phase grains is secured, and the effect of increasing _HcJ is easily obtained. Moreover, the decrease in B _r is also suppressed, and the desired magnetic properties are easily obtained.

In a more preferred structure of the sintered magnet of the present invention, a thin boundary phase is formed between the R'(Fe,Co) ₂ phase and the main phase. By separating the R'(Fe,Co) ₂ phase and the main phase by this thin boundary phase, the magnetic interaction between the two phases is weakened, and the room temperature _HcJ and the temperature dependence of _HcJ are further improved.
This boundary phase may be amorphous with a disordered atomic arrangement, or may have regular atomic arrangement. When the boundary phase is observed using a device such as a scanning transmission electron microscope (STEM), its composition contains 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. The content of R' is more preferably 25 atomic % or more, and even more preferably 30 atomic % or more. In addition to R', Fe, Co, and M, elements such as C, N, and O may be included.
The thickness of the boundary phase is preferably 0.1 nm or more and 20 nm or less. In this range, the effect of weakening the magnetic interaction between the R(Fe, Co) ₂ phase and the main phase is produced, and the reduction in the volume fraction of the main phase due to the formation of the boundary phase can be suppressed. The thickness is more preferably 0.2 nm or more and 10 nm or less, and particularly preferably 0.5 nm or more and 5 nm or less.

The Ce/R' of this thin boundary phase formed between the R'(Fe,Co) ₂ phase and the main phase is preferably higher than the Ce/R' of the grain boundary phase between two grains formed between the main phase grains. Since the boundary phase is adjacent to the R'(Fe,Co) ₂ 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, so that if the area of the main phase grain surface covered by this phase increases, the magnet will exhibit a higher room temperature _HcJ . The value of Ce/R' of the boundary phase is preferably 0.2 or more. It is more preferably 0.3 or more, and particularly preferably 0.35 or more.
In this way, by adopting a structural form in which a boundary phase with a high Ce/R' ratio is formed between the main phase grains and the R'(Fe, Co) _two phases, the magnetic interaction between the main phase and the R'(Fe, Co) _two phases is weakened, resulting in a high room temperature _HcJ and good temperature dependence of _HcJ .

The thickness of the boundary phase formed between the R'(Fe, Co) ₂ phase and the main phase and the grain boundary phase between two particles between the main phase grains can be calculated from the obtained HAADF (High-Angle Annular Dark Field) image by observing the areas where the main phase grains are adjacent to each other and the areas where the R'(Fe, Co) ₂ phase and the main phase are adjacent using, for example, an STEM device (JEM-ARM200F manufactured by JEOL Ltd.).

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., which are formed by the unavoidably mixed C, N, and O. From the viewpoint of suppressing deterioration of the magnetic properties, the volume ratio of these is preferably 10 volume % or less, and more preferably 5 volume % or less.

For example, the B-rich phase represented by R'1 _+ε (Fe,Co) _4B4 is preferably ₅ 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 the magnetic properties, it is preferable that the α-(Fe,Co) phase and the _R'2 (Fe,Co,M) ₁₇ phase are not included in the anisotropic rare earth sintered magnet of the present invention.

Next, the manufacturing method will be described. The anisotropic rare earth sintered magnet of the present invention is manufactured by powder metallurgy, but examples of methods for manufacturing a magnet having a structure in which the Ce/R' ratio differs between the center and outer shell of the main phase grains include the two-alloy method and the grain boundary diffusion method.

First, to prepare the raw alloy, metal raw materials, alloys, ferroalloys, etc. of R', Fe, Co, M are used, and the final sintered body is adjusted to have a predetermined composition after considering the raw material loss during the manufacturing process. These raw materials are melted in a high-frequency furnace or an arc furnace to prepare the alloy. The cooling from the molten metal may be performed by casting, or may be made into thin pieces by strip casting. In the case of the strip casting method, it is preferable to adjust the cooling rate to prepare the alloy so 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 are not sufficiently oriented in the process of molding in a magnetic field, resulting in a decrease in B _r . The average crystal grain size can be calculated, for example, by polishing the cross section of the alloy, etching it, and then observing the structure, drawing 20 lines parallel to the roll contact surface at equal intervals, and counting the intersections where these lines intersect with the grain boundary phase portion removed by etching. If α-Fe precipitates in the alloy, the alloy may be heat treated to remove the α-Fe and increase the amount of Nd ₂ Fe ₁₄ B type compound phase formed.

The above raw alloy is coarsely pulverized into powder with an average particle size of 0.05 to 3 mm by mechanical pulverization using a Braun mill or other means, or by hydrogen pulverization. Alternatively, the HDDR method (hydrogen disproportionation desorption recombination method) may be applied. The coarse powder is then finely pulverized using a ball mill or a jet mill using high-pressure nitrogen to obtain powder with an average particle size of 0.5 to 20 μm, more preferably 1 to 10 μm. Note that a lubricant may be added before or after the fine pulverization process, as necessary.

When the two-alloy method is used, two kinds of raw alloys with different compositions are prepared. Three or more kinds of alloys may be used. In this case, it is preferable to combine an alloy A mainly composed of a Nd ₂ Fe ₁₄ B type compound phase and having a relatively low Ce/R′ ratio with an alloy B having a relatively higher R′ composition ratio and Ce/R′ ratio than that, and adjust the average composition to a predetermined composition. These alloys are prepared by a casting method or a strip casting method, and then pulverized. The process of mixing the alloy powders may be performed in a coarse powder state before fine pulverization, or may be performed after fine pulverization.

Next, the alloy powder is molded into a green compact using a magnetic field press while aligning the magnetization easy axis of the alloy powder in an applied magnetic field. The molding is preferably performed in a vacuum, a nitrogen gas atmosphere, or an inert gas atmosphere such as Ar to suppress oxidation of the alloy powder.
The step of sintering the powder compact is carried out at a temperature of 800°C to 1200°C in a vacuum or inert atmosphere using a sintering furnace. Sintering is difficult to proceed at temperatures below 800°C, so a high sintered density cannot be obtained, and at temperatures above 1200°C, the main phase of the Nd ₂ Fe ₁₄ B type compound decomposes and α-Fe precipitates. The sintering temperature is 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. Sintering may be performed in a pattern in which the temperature is raised and then held at a constant temperature, or in order to refine the crystal grains, a two-stage sintering pattern in which the temperature is raised to a first sintering temperature and then held at a lower second sintering temperature for a predetermined time may be used. Sintering may also be performed multiple times, or a spark plasma sintering method or the like may be applied. The cooling rate after sintering is not particularly limited, but can be at least 600°C or less, preferably 200°C or less, preferably at a cooling rate of 1°C/min to 100°C/min, more preferably 5°C/min to 50°C/min. In order to improve the room temperature coercivity and the temperature characteristic of the coercivity, it is preferable to further perform 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 less, preferably 100°C or less, preferably 1°C/min to 100°C/min, more preferably 5°C/min to 50°C/min. The aging heat treatment may be performed multiple times. In addition, an intermediate heat treatment at 600 to 1000°C for 0.5 to 50 hours may be performed between the sintering heat treatment and the aging heat treatment.

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

By carrying out the above intermediate heat treatment and aging heat treatment under optimal conditions according to the composition, powder grain size, etc., an R-rich phase and an R(Fe, Co) ₂ phase are formed at the grain boundaries. In more preferable cases, an interparticle grain boundary phase is formed between adjacent main phase grains, and a thin boundary phase is formed between the R(Fe, Co) ₂ phase and the main phase grains. This leads to an increase in room temperature coercivity and an improvement in the temperature characteristics of the coercivity. The sintered body is cut and ground into a specified shape and magnetized to produce a sintered magnet.

1, in the sintered magnet by the two-alloy method, the main phase made _{of Nd2Fe14B} type compound is formed mainly by the components of alloy A, and the R'-rich phase, R'(Fe,Co) ₂ phase, and the outer shell of the main phase grain 10 are formed mainly by the components of alloy B. Therefore, the Ce/R' atomic ratio of the R'-rich phase and R'(Fe,Co) ₂ phase formed in the grain boundary portion 20 is higher than the Ce/R' atomic ratio inside the main phase grain. Also, some of the Ce in the grain boundary portion 20 replaces R' atoms in the surface layer portion of the main phase grain 10, forming a core-shell structure with different Ce concentrations in the center and outer shell.

On the other hand, in the grain boundary diffusion method, a sintered body is first produced by the single alloy method or the two-alloy method as described above. In this case, it is preferable that R' in the sintered body composition does not contain Ce.

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

The sintered body is heat-treated in a vacuum or in an inert gas atmosphere with Ce placed on the 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 the heat treatment is not particularly limited, but is preferably 1 to 20°C/min, and particularly preferably 2 to 10°C/min. The Ce placed on the sintered body penetrates into the inside of the sintered body via 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 grain 10 are replaced with Ce, and a core-shell structure with different Ce/R' ratios is formed between the center and outer shell portion of the main phase grain 10, and an R'-rich phase containing Ce or an R'(Fe, Co) ₂ phase is formed in the grain boundary portion 20, and _HcJ increases.

As with the two-alloy method, it is preferable to further subject the diffusion heat-treated sintered body to aging heat treatment at 300 to 800°C for 0.5 to 50 hours in order to improve the room temperature coercivity and the temperature characteristics of the coercivity.

In order to form a thin boundary phase between the main phase grains and the R(Fe, Co) ₂ grain boundary phase, the sintered body after the diffusion treatment may be subjected to an intermediate heat treatment similar to that in the two-alloy method, but in this case, it can 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 an R(Fe, Co) ₂ phase are formed at the grain boundary portion, and further a thin boundary phase is formed between the R(Fe, Co) ₂ phase and the main phase grains. In a more preferable case, a two-particle grain boundary phase is formed between adjacent main phase grains, and the room temperature coercivity and the temperature characteristic of the coercivity are increased.

To further improve the magnetic properties, Dy or Tb may be placed on the surface of the sintered body, either separately or simultaneously with Ce, and a diffusion heat treatment may be performed.

The anisotropic rare earth sintered magnet of the present invention thus produced exhibits a residual magnetic flux density B _r of at least 12 kG or more at room temperature, and a coercive force H _cJ of at least 10 kOe. The temperature coefficient β of the coercive force exhibits the following characteristic: β≧(0.01×H _{cJ(room temperature)} −0.720)%/K, where β=ΔH _cJ /ΔT×100/H _{cJ(room temperature)} , (ΔH _cJ =H _{cJ(room temperature)} −H _{cJ(140° C.)} , ΔT=room temperature−140(° C.)). It is even more preferable if β≧(0.01×H _{cJ(room temperature)} −0.7)%/K. The anisotropic rare earth sintered magnet of the present invention exhibits a smaller change in coercive force with temperature than a Ce-free Nd—Fe—B sintered magnet, making it suitable for use at high temperatures.

The present invention will be specifically explained below with examples and comparative examples, but the present invention is not limited to the following examples.

[Example 1]
Using Nd metal, Pr metal, electrolytic iron, Co metal, ferroboron, Al metal, and Cu metal, the composition was adjusted to Nd 10.6 atomic %, Pr 2.7 atomic %, Co 1.0 atomic %, B 6.0 atomic %, Al 0.5 atomic %, Cu 0.1 atomic %, and the balance Fe, and 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 to produce an alloy ribbon with a thickness of about 0.2 to 0.4 mm. The cross section of this alloy was polished and etched, and then the structure was observed with a laser microscope (Olympus Corporation, LEXT OLS4000). The observation points were located at a position about 0.15 mm from the surface where the ribbon contacted the cooling roll, and 20 observation points were observed. 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 portion removed by etching were counted to calculate the average grain boundary phase interval, which was 4.7 μm. This alloy was subjected to hydrogen absorption treatment at room temperature, and then subjected to dehydrogenation treatment by heating in a vacuum at 400 ° C. to obtain a coarse powder (this is referred to as the actual 1A powder). Next, an alloy ingot was manufactured using Ce metal and electrolytic iron as raw materials, with a composition adjusted to 33 atomic % Ce and the remainder Fe using a high-frequency induction furnace, and was heat-treated at 870 ° C. for 20 hours, and then mechanically pulverized to obtain a coarse powder (referred to as the actual 1B powder). The actual 1A powder and the actual 1B powder were mixed in a weight ratio of 93:7, and then pulverized with a jet mill in a nitrogen stream to obtain a fine powder with an average particle size of 3.1 μm. Next, the fine powder was filled into a mold of a molding machine in an inert gas atmosphere, and while being oriented in a magnetic field of 15 kOe (=1.19 MA/m), it was pressure-molded in a direction perpendicular to the magnetic field at a pressure of 0.6 Ton/ ^cm2 . The obtained powder compact was sintered in a vacuum at 1040°C for 3 hours, then cooled to room temperature, taken out, and further subjected to a heat treatment at 510°C for 2 hours to obtain a sintered body sample of Example 1.

The obtained sintered body sample was analyzed by inductively coupled plasma optical emission spectroscopy (ICP-OES) using a high-frequency inductively coupled plasma optical emission spectroscopy analyzer (Hitachi High-Tech Science Corporation, SPS3520UV-DD), and the composition was Nd _9.9 Pr _2.5 Ce _1.8 Fe _bal. Co _1.0 B _5.6 Al _0.5 Cu _0.1 . 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 B type. When the structure of the sintered body and the composition analysis of each phase were performed using an EPMA device (JXA-8500F, manufactured by JEOL Ltd.), a core / shell structure with different compositions was formed in the center and outer shell of the main phase grain. Ce was not contained in R' in the center corresponding to the core, and Ce was contained in R' in the outer shell of the grain. It was also confirmed that the grain boundary portion contained 1 volume % or more of each of the R'-rich phase and the R'(Fe, Co) ₂ phase. The volume ratio of each phase was calculated as being equal to the area ratio in the backscattered electron image. The α-Fe phase and the R' ₂ (Fe, Co, M) ₁₇ phase were not observed. The total phase ratio was less than 100% because oxides and other phases were also present. Based on the analysis value of the R'(Fe, Co) ₂ phase, an alloy of the same composition was prepared by arc melting and homogenized at 800°C for 10 hours, and then magnetization-temperature measurements were performed using a VSM, revealing that the Curie temperature _Tc was 66°C.

The average crystal grain size of the main phase, calculated as described above from the results of etching and observing the sintered body sample, was 4.3 μm. When the magnetic properties were measured with a B-H tracer, the values of B _r 14.0 kG and H _cJ 13.6 kOe were obtained at room temperature. The temperature coefficient β of H _cJ was -0.575%/K. Table 1 shows the ICP compositional analysis values of the sintered body, the average crystal grain size, and the crystal structure of the main phase. Table 2 shows the conditions of the sintering heat treatment and the aging heat treatment, and the results of the magnetic properties measured with the B-H tracer, and Table 3 shows the compositional analysis values of each phase measured with an EPMA.

[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 prepare a strip cast alloy ribbon. The average grain boundary phase interval calculated from the cross-sectional image of this alloy was 4.4 μm. This alloy was subjected to hydrogen absorption treatment and dehydrogenation treatment by heating at 400 ° C in a vacuum to obtain a coarse powder, and then pulverized with a jet mill in a nitrogen gas flow to obtain a fine powder with an average particle size of 3.1 μm. After being made into a powder compact by pressure molding in a magnetic field, it was sintered in a vacuum at 1040 ° C for 3 hours, cooled to room temperature, taken out, and further subjected to heat treatment at 510 ° C for 2 hours to obtain a sintered body sample of Comparative Example 1.

From the 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 . The main phase was confirmed to have a Nd ₂ Fe ₁₄ B type crystal structure by X-ray diffraction. When the structure observation and composition analysis of each phase were performed with an EPMA device, the composition in the main phase grain was almost uniform, and there was no difference in Ce concentration between the center and the outer shell. In addition, an R'-rich phase was present in the grain boundary, but the R'(Fe, Co) ₂ phase could not be confirmed. The average crystal grain size of the main phase was 4.0 μm. The magnetic properties measured with a B-H tracer were B _r 13.7 kG, 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.

[Example 2, Comparative Example 2]
In Example 2, similarly to Example 1, a strip cast alloy ribbon having a composition of Nd 12.8 atomic %, Co 1.0 atomic %, B 5.9 atomic %, Al 0.2 atomic %, Zr 0.05 atomic %, and the balance Fe, with a thickness of about 0.2 to 0.4 mm and an average grain boundary phase interval of 3.9 μm was prepared, and subjected to hydrogen absorption treatment and dehydrogenation treatment to obtain a coarse powder (actual 2A powder). On the other hand, an alloy adjusted to have a composition of Ce 80 atomic %, Cu 10 atomic %, and the balance 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 quenched alloy ribbon having a thickness of about 100 to 250 μm. This alloy ribbon was crushed into a coarse powder (actual 2B powder) by ball milling. The powders of Example 2A and Example 2B were mixed in a weight ratio of 96:4, and then pulverized in a jet mill in a nitrogen gas flow to obtain fine powder with an average particle size of 2.8 μm. The powder was then pressed in a magnetic field to obtain a green compact, which was then sintered in a vacuum at 1020° C. for 2 hours, cooled to room temperature, removed, and further heat-treated at 530° C. for 4 hours to obtain a sintered sample of Example 2.

In Comparative Example 2, a strip cast alloy ribbon with a composition of Nd 7.8 atomic %, Ce 5.0 atomic %, Co 1.0 atomic %, B 5.9 atomic %, Al 0.2 atomic %, Zr 0.05 atomic %, and the remainder Fe, was prepared with a thickness of about 0.2 to 0.4 mm and an average grain boundary phase spacing of 4.2 μm, and was subjected to hydrogen absorption and dehydrogenation treatment to obtain a coarse powder (ratio 2A powder). On the other hand, an alloy adjusted to a composition of Nd 80 atomic %, Cu 10 atomic %, and the remainder 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 quenched alloy ribbon with a thickness of about 100 to 250 μm. This alloy ribbon was crushed into a coarse powder (ratio 2B powder) by ball milling. The powders of ratio 2A and ratio 2B were mixed in a weight ratio of 96:4, then pulverized in a jet mill in a nitrogen stream to obtain fine powder with an average particle size of 2.8 μm. After being pressed in a magnetic field to produce a powder compact, the compact was sintered in a vacuum at 1020°C for 2 hours, cooled to room temperature, removed, and further heat-treated at 530°C for 4 hours to obtain a sintered sample of Comparative Example 2.

The compositions of the sintered bodies of Example 2 and Comparative Example 2 were Nd _12.4 Ce _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 , respectively, as determined by ICP analysis. When the structure was observed, it was found that in Example 2, the center did not contain Ce, and many main phase grains containing Ce were present in the outer shell of the grains, and the grain boundaries contained 1 volume % or more of R'-rich phase and R'(Fe, Co) ₂ phase. The _Tc of the alloy of the same composition produced by arc melting based on the analysis value of R'(Fe, Co) ₂ phase was 74°C. On the other hand, in Comparative Example 2, both the center and outer shell of the main phase grains contained Ce, and the Ce/R' ratio was higher in the center of the grain than in the outer shell of the grain. In addition, R'(Fe,Co) ₂ phase and R'Cu ₂ phase were formed in the grain boundary, and no R'-rich phase was observed. 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 better than Comparative Example 2 in both room temperature magnetic properties and _HcJ temperature properties.

[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, were prepared by arc melting, and mixed in a weight ratio of 94:6 as coarse powders as in Example 1. A powder compact produced by jet mill pulverization and pressure molding in a magnetic field was sintered in a vacuum at 1010°C for 3 hours. Then, an aging heat treatment was performed at 480°C for 1 hour to obtain a sintered sample.

In Example 4, a strip cast alloy with a composition of Nd 12.8 atomic %, B 6.0 atomic %, Al 0.5 atomic %, Cr 0.2 atomic %, Ti 0.3 atomic %, and the balance Fe was prepared, and a cast alloy with a composition of Ce 28 atomic %, Gd 7 atomic %, Co 30 atomic %, and the balance Fe was prepared, and mixed in a weight ratio of 90:10 as coarse powder as in Example 1. The powder compact produced by jet mill crushing and pressure molding 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 less at a cooling rate of 3.8°C/min, and then aged at 600°C for 3 hours to obtain a sintered body sample.

In Example 5, a strip cast alloy adjusted to have a composition of 13.0 atomic % Nd, 6.0 atomic % B, and the balance Fe, and an arc-melted alloy adjusted to have a composition of 56 atomic % Ce, 9 atomic % Y, 10 atomic % Si, 8 atomic % Ga, and the balance Co were prepared and mixed in a weight ratio of 95:5 as coarse powders as in Example 1. The powder compacts prepared by jet mill crushing and pressure molding in a magnetic field were sintered in a vacuum at 1060°C for 2 hours. The obtained sintered bodies were heat-treated at 960°C for 2 hours, cooled to below 500°C at a cooling rate of 4.5°C/min, and then aged at 680°C for 3 hours to obtain sintered body samples.

The results of Examples 3 to 5 are shown in Tables 1 to 3. In all of the sintered body structures, the grain center did not contain Ce, and many main phase grains containing Ce were present in the grain outer shell, and the grain boundary portion contained an R'-rich phase and an R'(Fe, Co) ₂ phase in a total amount of 1 volume % or more. In addition, the magnetic properties were good, with _HcJ at room temperature being 10 kOe or more and the temperature coefficient β of _HcJ being (0.01 x HcJ _{(room temperature)} - 0.720) %/K or more.

[Example 6, Comparative Example 3]
The composition was adjusted using Nd metal, electrolytic iron, Co metal, ferroboron, and Al metal to produce a strip cast alloy ribbon. The average grain boundary phase spacing calculated from the cross-sectional image of this alloy was 4.8 μm. This alloy was subjected to hydrogen absorption treatment and dehydrogenation treatment by heating at 400 ° C in a vacuum to obtain a coarse powder, which was then pulverized with a jet mill in a nitrogen gas flow to obtain a fine powder with an average particle size of 3.5 μm. A powder compact was formed by pressure molding in a magnetic field, and sintered in a vacuum at 1040 ° C for 3 hours. The obtained sintered body was then cut to a size of 10 × 10 × 3 mm.
Next, an alloy ingot was manufactured using Ce metal, Dy metal, electrolytic iron, Co metal, and Cu metal as raw materials, and the composition was adjusted to Ce 25 atomic %, Dy 8 atomic %, Co 30 atomic %, Cu 10 atomic %, and the balance Fe using a high-frequency induction furnace. After heat treatment at 420 ° C for 20 hours, it was pulverized in a ball mill to obtain a powder with an average particle size of 14.6 μm. The above sintered body was immersed in a liquid in which this powder and ethanol were mixed and stirred at a weight ratio of 1:1, pulled up, dried by blowing air, and powder coating was performed on 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 less 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 sintered body sample of Comparative Example 3 was prepared by subjecting the above powder coating and diffusion heat treatment to only aging heat treatment at 560° C. for 2 hours in an Ar gas atmosphere.

From the ICP analysis, the compositions of the sintered bodies 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 and Nd _14.0 Fe _bal. Co _0.4 B _6.0 Al _0.1 , respectively. From the EPMA structure observation at a depth of 500 μm from the surface of the sintered body, in Example 6, the center did not contain Ce, and many main phase grains containing Ce were present in the outer shell of the grains, and the R'-rich phase and the R'(Fe, Co) ₂ phase were each present at 1 vol % or more in the grain boundary. The _Tc of the alloy of the same composition produced by arc melting based on the analysis value of the R'(Fe, Co) ₂ phase was 131 ° C. On the other hand, Comparative Example 3 did not contain Ce, and an R'-rich phase was present in the grain boundary, but no R'(Fe, Co) ₂ phase was observed. 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 _HcJ temperature characteristics than Comparative Example 3.

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

In Example 8, a strip cast alloy was prepared with a composition adjusted to Nd 14.1 atomic %, B 6.0 atomic %, Al 0.5 atomic %, Cu 0.1 atomic %, and the balance Fe, and a strip cast alloy ribbon was prepared. The average grain boundary phase spacing calculated from the cross-sectional image of this alloy was 4.8 μm. This alloy was subjected to hydrogen absorption treatment and dehydrogenation treatment by heating at 400 ° C in a vacuum to obtain a coarse powder, which was then pulverized with a jet mill in a nitrogen gas flow to obtain a fine powder with an average particle size of 3.3 μm. A powder compact was formed by pressure molding in a magnetic field, and sintered in a vacuum at 1030 ° C for 2 hours. The obtained sintered body was cut to a size of 10 × 10 × 3 mm.
Next, the sintered body was immersed in a liquid in which Ce oxide powder and pure water were mixed and stirred in a weight ratio of 3:2, pulled out, dried by blowing air, and the powder was applied to the surface of the sintered body. This sample was subjected to a diffusion heat treatment at 880°C in a vacuum for 20 hours, cooled to 450°C or less 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.

In Example 9, a strip cast alloy adjusted to have a composition of Nd 14.5 atomic %, Co 1.0 atomic %, B 6.2 atomic %, Al 0.2 atomic %, Cu 0.1 atomic %, Zr 0.05 atomic %, and the balance Fe, and an alloy adjusted to have a composition of Ce 30 atomic %, Co 35 atomic %, and the balance Fe and arc melted were prepared, and mixed as in Example 1 in a weight ratio of 95:5 as coarse powder, and pulverized in a jet mill in a nitrogen gas flow to obtain a fine powder with an average particle size of 3.7 μm. A powder compact was formed by pressure molding in a magnetic field, and sintered in a vacuum at 1020 ° C for 3 hours. The obtained sintered body was cut to a size of 10 × 10 × 3 mm.
Next, the above sintered body was immersed in a liquid in which Tb oxide powder and pure water were mixed and stirred in a weight ratio of 1:1, pulled out, dried by blowing air, and the powder was applied to the surface of the sintered body. This sample was subjected to a diffusion heat treatment at 830°C in a vacuum for 20 hours, cooled to 500°C or less at a cooling rate of 5°C/min, and further subjected to an aging heat treatment at 530°C in an Ar gas atmosphere for 1.5 hours to obtain a sintered body sample of Example 9.

The results of Examples 7 to 9 are shown in Tables 1, 2, and 4. In all of the sintered body structures, the center did not contain Ce, and many main phase grains containing Ce were present in the outer shell of the grains, and an R'-rich phase and an R'(Fe, Co) ₂ phase were each present at 1 volume % or more in the grain boundaries. In addition, the magnetic properties were good, with _HcJ at room temperature being 10 kOe or more and the temperature coefficient β of _HcJ being (0.01 x HcJ _{(room temperature)} - 0.720) %/K or more.

[Example 10, Comparative Example 4]
A strip cast alloy ribbon having a composition of Nd 13.5 atomic %, B 6.0 atomic %, Al 0.5 atomic %, Cu 0.2 atomic %, and the balance Fe, with 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 hydrogen absorption treatment and dehydrogenation treatment to obtain a coarse powder (actual 10A powder). Next, an alloy having a composition adjusted to Ce 35 atomic %, Co 10 atomic %, and the balance Fe was produced using an arc melting furnace, and after heat treatment at 850 ° C. for 15 hours, it was mechanically pulverized to obtain a coarse powder (actual 10B powder). The actual 10A powder and the actual 10B powder were mixed in a weight ratio of 92:8, and then pulverized with a jet mill in a nitrogen stream to obtain a fine powder with an average particle size of 3.6 μm. After being pressure-molded in a magnetic field to produce a green compact, the green compact was sintered in a vacuum at 1000°C for 2 hours, cooled to room temperature, removed, and further heat-treated at 500°C for 3 hours to obtain a sintered compact sample of Example 10.
On the other hand, a sample prepared in the same manner as in Example 10 up to the sintering step was subjected to a heat treatment at 980° C. for 1 hour and then cooled in an Ar atmosphere to prepare Comparative Example 4.

From the ICP analysis, the sintered body composition of Example 10 and Comparative Example 4 was Nd _12.5 Ce _2.1 Fe _bal. Co _0.7 B _5.8 Al _0.4 Cu _0.1 . From the EPMA structure observation, both did not contain Ce in the center, and many main phase grains containing Ce were present in the outer shell of the grain. In Example 10, the R'-rich phase and the R'(Fe, Co) ₂ phase were present at 1 volume % or more in the grain boundary. The T _c of the alloy of the same composition prepared by arc melting based on the analysis value of the R'(Fe, Co) ₂ phase was 70 ° C. On the other hand, in Comparative Example 4, the R'-rich phase was present in the grain boundary, but the R'(Fe, Co) ₂ phase could not be confirmed. The average crystal grain size of the main phase was 4.9 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 _HcJ than Comparative Example 4, and the temperature characteristics of _HcJ were also good.

[Example 11]
A strip cast alloy ribbon having a composition of Nd13.5 atomic %, B5.9 atomic %, Co1.0 atomic %, Al0.5 atomic %, Cu0.2 atomic %, Zr0.1 atomic %, and the balance Fe, with a thickness of about 0.2 to 0.4 mm and an average grain boundary phase spacing of 4.2 μm was produced, and subjected to hydrogen absorption treatment and dehydrogenation treatment to obtain a coarse powder (actual 11A powder). Next, an alloy ingot adjusted to have a composition of Ce33.3 atomic %, Co1.0 atomic %, and the balance Fe was produced using an arc melting furnace, and after heat treatment at 860 ° C. for 18 hours, it was mechanically pulverized to obtain a coarse powder (actual 11B powder). The actual 11A powder and the actual 11B powder were mixed in a weight ratio of 93:7, and then pulverized with a jet mill in a nitrogen stream to obtain a fine powder with an average particle size of 2.9 μm. After forming a powder compact by pressure molding in a magnetic field, the compact was sintered in a vacuum at 1020°C for 3 hours, cooled to room temperature, and removed. Next, intermediate heat treatment was performed at 900°C for 1 hour in an Ar atmosphere, and then cooled to 450°C or less at a cooling rate of 5°C/min, followed by low-temperature heat treatment at 510°C for 3 hours to obtain a sintered compact sample of Example 11.

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 _0.1 by ICP analysis. EPMA structure observation showed that there were many main phase grains that did not contain Ce in the center and contained Ce in the outer shell of the grains. In addition, R'-rich phase and R'(Fe, Co) ₂ phase were present at 1 volume % or more each in the grain boundary. The T _c of an alloy of the same composition prepared by arc melting based on the analysis value of R'(Fe, Co) ₂ phase was 68 ° C. The average crystal grain size of the main phase was 3.9 μm. The results are shown in Tables 1, 2 and 5.

Using a FIB-SEM device (Scios manufactured by FEI), a specimen for observation was cut out from the sample of Example 11, and observed with a STEM device (JEM-ARM200F manufactured by JEOL Ltd.). As shown in the HAADF image in FIG. 5, it was confirmed that a boundary phase was formed between the R _' (Fe _, Co _{) 2} phase at the grain boundary and the main phase. The thickness of this boundary phase was 1.4 nm on average, and the composition of the boundary phase measured by EDS analysis _{was Nd22.5Ce13.5Febal.Co3.0Cu1.7} . On the other hand, the EDS analysis composition of the adjacent _R '(Fe,Co _{) 2} phase was _{Nd14.7Ce19.5Febal.Co2.3Cu0.1 .} This shows that this boundary phase is a phase having a different composition from the R'(Fe _, Co) ₂ phase.
In another location of the same sample, a grain boundary phase with an average thickness of about 2.4 nm exists between adjacent main phase grains, and its average composition was Nd _26.8 Ce _6.9 Fe _bal. Co _7.4 Cu _12.5 Zr _0.5 as determined by EDS analysis. From this, the Ce/R' ratios calculated for the boundary phase formed between the main phase and the R'(Fe,Co) ₂ phase, and the grain boundary phase between two grains between the main phase grains are 0.37 and 0.20, respectively, indicating that the former exhibits a higher Ce/R' ratio.

[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 phase spacing of 4.0 μm was produced, which was subjected to hydrogen absorption and dehydrogenation treatments, and then pulverized in a jet mill in a nitrogen stream to obtain a fine powder with an average particle size of 3.0 μm. A powder compact was formed by pressure molding in a magnetic field, and the sintered body obtained by sintering in a vacuum at 1040° C. for 2 hours was cut to a size of 10 × 10 × 3 mm.
Next, a Ce film was formed on one surface of the 10×10 mm surface of the sintered body by sputtering for 90 minutes using a target having a composition of Ce ₃₀ Fe _bal. Co ₂₀ Al ₂₀ Cu ₅ V ₅ , a diameter of 2 inches, and a thickness of 3 mm, at an input power of 250 W and an Ar pressure of 0.4 Pa. This sample was subjected to a diffusion heat treatment at 840° C. in a vacuum for 25 hours, cooled to 500° C. or less at a cooling rate of 4.5° C./min, and then subjected to an aging heat treatment at 540° C. in an Ar gas atmosphere for 3 hours to obtain a sintered body sample of Example 12.

The composition of the sintered body 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 by ICP analysis. EPMA structure observation showed that the main phase grains contained no Ce in the center and many Ce-containing grains in the outer shell. In addition, the grain boundary part contained 1 vol% or more of R'-rich phase and R'(Fe,Co) ₂ phase. The Tc of the _alloy of the same composition produced by arc melting based on the analysis value of R'(Fe,Co) ₂ phase was 78°C. The results are shown in Tables 1, 2 and 5.

The structure of Example 12 was observed by STEM, and it was confirmed that 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 formed between the R'(Fe, Co) 2 phase and the main phase. From this, the Ce/R' of the boundary phase was calculated to be 0.38. Meanwhile, in another part of the same sample, a grain boundary phase with an average thickness of about 1.8 nm was present 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) This shows that the Ce/R' of the boundary phase formed between the main phase and the R'(Fe, Co) _two phases is higher than the Ce/R' of the grain boundary phase between the two particles.

11 Main phase (region with high Ce/R′)
12 Main phase (low Ce/R′ region)
21 R'-rich phase 22 R'(Fe, Co) ₂ phase 31 Intergranular grain boundary phase formed between adjacent main phase grains 32 Boundary phase formed between R'(Fe, Co) ₂ phase and main phase

Claims (18)

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 (R is two or more elements selected from rare earth elements, with Nd and Ce being essential, 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, and 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), and the main phase is Nd ₂ Fe ₁₄ An anisotropic rare earth sintered magnet is characterized in that it is made of a B-type crystal compound, has main phase grains in which the Ce/R' ratio (R' is one or more elements selected from rare earth elements, and Nd is an essential element) at the center of the grain is lower than the Ce/R' ratio at the outer periphery of the grain, and has an R'-rich phase containing Ce and an R'(Fe, Co) ₂ phase containing Ce at grain boundaries. _2. The anisotropic rare earth sintered magnet according to claim 1, characterized in that 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)2 phase. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that there are main phase grains in which the central R' does not contain Ce. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the main phase grains have a central R' of Nd or Nd and Pr. 3. The anisotropic rare earth sintered magnet according to claim 1, wherein the R'(Fe, Co) ₂ phase is a phase that exhibits ferromagnetism or ferrimagnetism at room temperature or above. 3. The anisotropic rare earth sintered magnet according to claim 1, wherein the Ce/R' ratio in the R'(Fe, Co) ₂ phase is higher than the Ce/R' ratio in the outer shell portion of the main phase grains. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the Ce/R' ratio in the R'-rich phase is higher than the Ce/R' ratio in the outer shell of the main phase grains. 3. The anisotropic rare earth sintered magnet according to claim 1, wherein the R'-rich phase and the R'(Fe, Co) ₂ phase are contained in a total amount of 1% by volume or more. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the Ce/R' ratio in the composition of the sintered body is 0.01 or more and 0.3 or less. The anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that the B-rich phase contained in the sintered magnet is 5% by volume or less. An anisotropic rare earth sintered magnet according to claim 1 or 2, characterized in that a grain boundary phase is formed between adjacent main phase grains. 12. The anisotropic rare earth sintered magnet according to claim ₁₁ , characterized in that the Ce/R' in the boundary phase formed between the main phase and the R'(Fe, Co) two phases is higher than the Ce/R' in the intergranular grain phase formed between adjacent main phase grains. 3. The anisotropic rare earth sintered magnet according to claim 1, characterized in that the coercive force HcJ _(room) at room temperature is 10 kOe or more, and the value of the temperature coefficient of coercive force β is expressed as β≧(0.01×HcJ _{(room temperature)} −0.720)%/K. The method for producing an anisotropic rare earth sintered magnet according to claim _{1 or} 2, characterized in that an alloy containing a _Nd2Fe14B type crystal compound phase and an alloy having a higher R' composition ratio and Ce/R' ratio than that of the alloy are crushed and mixed, compacted into a green body by applying a magnetic field, and then sintered at a temperature of 800°C to 1200°C. 3. The method for producing an anisotropic rare earth sintered magnet according to claim 1, further comprising the steps of: pulverizing an alloy containing a compound phase _{of Nd2Fe14B} type crystals; compacting the powder in an applied magnetic field to form a green body; sintering the green body at a temperature of 800°C to 1200°C; contacting the sintered body with a material containing Ce; and heat treating the sintered body at a temperature of 600°C to the sintering temperature to diffuse Ce into the sintered body. The method for producing an anisotropic rare earth sintered magnet according to claim 15, characterized in that the Ce-containing material brought into contact with the sintered body is one or more selected from Ce metal, a Ce-containing alloy, and a Ce-containing compound, and is in the form of one or more selected from powder, thin film, ribbon, foil, and gas. The method for producing an anisotropic rare earth sintered magnet according to claim 14, characterized in that the sintered body is subjected to a heat treatment at a temperature of 300 to 800°C. The method for producing anisotropic rare earth sintered magnets according to claim 14, characterized in that the sintered body is heat-treated at a temperature of 600 to 1000°C, then cooled to at least 550°C or less at a cooling rate of 1°C/min to 50°C/min, and further heat-treated at a temperature of 300 to 800°C.