Anisotropic rare earth sintered magnet and method for producing the same

Abstract

To provide an anisotropic rare earth sintered magnet having a ThMn12 type crystal compound as the main phase, which exhibits good magnetic properties, and a method for manufacturing the same.SOLUTION: An anisotropic rare earth sintered magnet contains 80% by volume or more of a main phase consisting of a ThMn12 type crystal compound, the average crystal grain size of the main phase is 1 μm or more, an intergrain boundary phase 13 is formed between adjacent main phase grains 11 and 12. A manufacturing method of the anisotropic rare earth sintered magnet includes grinding an alloy containing a compound phase of the ThMn12 type crystal compound, and performing sintering at a temperature of 800°C or higher and 1400°C or lower after compacting into a compact while applying a magnetic field.SELECTED DRAWING: Figure 1

Description

[Technical field]

[0001] The present invention relates to a method for producing ThMn 12 The present invention relates to an anisotropic rare earth sintered magnet having a main phase of a compound of a type crystal, and a method for producing the same. [Background technology]

[0002] Rare earth magnets, especially Nd-Fe-B sintered magnets, are becoming increasingly popular due to the electrification of automobiles and Due to the trend towards higher performance and lower power consumption in industrial motors, demand is expected to grow and production volume is expected to increase further. However, there are concerns that the supply and demand balance of rare earth raw materials may be disrupted in the future, and in recent years, research into reducing the amount of rare earth used in rare earth magnets has been attracting attention. 12 The compound with the R type crystal structure is 2 Fe 14 Since it has a lower rare earth content than B compounds and has good magnetic properties, it is being actively researched as a next-generation magnet material.

[0003] For example, in Patent Document 1, ThMn 12 A permanent magnet made of an alloy containing a hard magnetic phase and a nonmagnetic phase with a tetragonal crystal structure has been reported. It is shown that adding at least one element selected from Cu, Si, Mg, Sn, Pb, and In to an intermetallic compound mainly consisting of rare earth elements and Fe precipitates a phase that has a lower melting point than the main phase and is nonmagnetic.

[0004] In addition, in Patent Document 2, a main phase and a grain boundary phase are included, and the main phase is ThMn 12An R-T compound having a type crystal structure (R is one or more rare earth elements including La as an essential element, T is Fe, or Fe and Co, or a part thereof is replaced by M (one or more selected from Ti, V, Cr, Mo, W, Zr, Hf, Nb, Ta, Al, Si, Cu, Zn, Ga, and Ge)), and the grain boundary phase has a cubic crystal structure. A rare earth permanent magnet having a La-rich phase σ with a La composition ratio of 20 at% or more and a cross-sectional area ratio of 20% or more has been reported. By including a non-magnetic cubic La-rich phase at the grain boundary, a magnetic separation effect between the main phases and a reduction effect of the interfacial strain between the grain boundary phase and the main phase are said to be obtained.

[0005] In Patent Document 3, a main phase having a ThMn 12 type crystal structure and a sub-phase containing any one of Sm 5 Fe 17 system phase, SmCo 5 system phase, Sm 2 O 3 system phase, and Sm 7 Cu 3 system phase have been reported. The volume fraction of the sub-phase is 2.3 to 9.5%. Among these sub-phases, the Sm 5 Fe 17 system phase and the SmCo 5 system phase are magnetic phases showing higher magnetic anisotropy than the main phase. By isolating each crystal grain of the main phase and preventing the movement of magnetic walls in the main phase, the magnetization and coercivity of the magnet are improved. On the other hand, the Sm 2 O 3 system phase and the Sm 7 Cu 3 system phase are non-magnetic phases. By isolating each crystal grain of the main phase, it is said that the propagation of magnetization reversal of the main phase to the surroundings is prevented, and the magnetization and coercivity of the magnet are improved. Also, the Sm 7 Cu 3 system phase is described as a non-equilibrium phase.

[0006] In Patent Document 4, it has a main phase and one or more sub-phases, and the composition of the entire alloy is R(Fe,Co) w-z Ti z Cu αAn alloy for rare earth magnets satisfying (R is at least one rare earth element, 8 ≦ w ≦ 13, 0.42 ≦ z < 0.70, 0.40 ≦ α ≦ 0.70) has been reported. Also, it is described that the secondary phase is mainly a crystal phase with a Cu composition of 50 mol% or more of the entire secondary phase, and the crystal structure of the secondary phase is of the KHg 2 type.

[0007] In Patent Document 5, R x Fe 100-x-y (V 1-a Si a ) y (R is one or more rare earth elements including Y, x = 5.5 to 18 atomic%, y = 8 to 20 atomic%, a = 0.05 to 0.7), and rare earth permanent magnets having a main phase with a ThMn 12 type body-centered tetragonal crystal structure have been reported. This composition alloy consists of a main phase and a rare earth-rich phase, and it is described that it does not contain an RFe 2 phase.

Prior Art Documents

Patent Documents

[0008]

Patent Document 1

Patent Document 2

Patent Document 3

Patent Document 4

Patent Document 5

Summary of the Invention

Problems to be Solved by the Invention

[0009] As described above, ThMn 12In order to obtain good magnetic properties in a magnet having a [[ID=]] type compound as the main phase, it has been proposed to have a structure composed of a main phase and a grain boundary phase, similar to Nd-Fe-B magnets, and non-magnetic phases such as La-rich phases (Patent Document 2) and R-Cu phases (Patent Documents 1 and 4) have been studied as the grain boundary phase. However, in reality, these phases tend to segregate at grain boundary triple points and it is difficult to form an inter-particle grain boundary phase, and there has been a problem that it is difficult to obtain a structure in which the surface of the main phase grains is covered by the grain boundary phase.

[0010] In addition, in Patent Document 3, Sm 5 Fe 17 -based phases and SmCo 5 -based phases, which are magnetic phases showing high magnetic anisotropy, surround the surface of the main phase grains, and the coercivity is improved by pinning magnetic domain walls with this phase. However, it is difficult to realize a tissue morphology in which the surface of the crystal grains of the ThMn 12 -type compound is surrounded by Sm 5 Fe 17 -based phases and SmCo 5 -based phases.

[0011] On the other hand, in Patent Document 5, an alloy composed of a ThMn 12 main phase and an R-rich phase is proposed. However, in reality, since the composition range in which only two phases are formed in the R-Fe-V-Si quaternary system is extremely limited, it is difficult to reproduce this structure with good reproducibility.

[0012] The present invention has been made in view of the above problems, and an object thereof is to provide an anisotropic rare earth sintered magnet having a ThMn 12 -type crystal compound as the main phase.

Means for Solving the Problems

[0013] As a result of intensive studies to achieve the above object, the present inventors have found that in an anisotropic rare earth sintered magnet having a ThMn 12 -type crystal compound as the main phase, a high coercivity is exhibited when an inter-particle grain boundary phase is formed between adjacent main phase grains, and the present invention has been completed.

[0014] 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 1-a Zr a ) x (Fe 1-b Co b ) 100-x-y (M 1 1-c M 2 c ) y (where R is one or more selected from rare earth elements and Sm is essential, M 1 is one or more elements selected from the group consisting of V, Cr, Mn, Ni, Cu, Zn, Ga, Al, Si, M 2 is one or more elements selected from the group consisting of Ti, Nb, Mo, Hf, Ta, W, and x, y, a, b, c are 7 ≦ x ≦ 15 at%, 4 ≦ y ≦ 20 at%, 0 ≦ a ≦ 0.2, 0 ≦ b ≦ 0.5, 0 ≦ c ≦ 0.9, respectively), and the anisotropic rare earth sintered magnet contains 80% by volume or more of a main phase composed of a ThMn 12 -type crystal compound, the average crystal grain size of the main phase is 1 μm or more, and an intergranular phase between adjacent main phase grains is formed. (2) The anisotropic rare earth sintered magnet according to (1), wherein the intergranular phase between the two particles contains 20% by atom or more of R. (3) The anisotropic rare earth sintered magnet according to (1) or (2), wherein the intergranular phase between the two particles has a thickness of 0.5 nm or more. (4) The anisotropic rare earth sintered magnet according to any one of (1) to (3), characterized in that it contains an R-rich phase at the grain boundary. (5) The anisotropic rare earth sintered magnet according to any one of (1) to (4), characterized in that it contains an R(Fe,Co) 2 phase at the grain boundary. (6) The anisotropic rare earth sintered magnet according to (4) or (5), characterized in that the total of the R-rich phase and the R(Fe,Co) 2 phase is 1% by volume or more. (7) The Sm / R ratio inside the main phase grains is 2The anisotropic rare earth sintered magnet according to any one of (4) to (6), characterized in that the Sm / R ratio of the phase is lower than that of the matrix phase. (8) The anisotropic rare earth sintered magnet according to any one of (1) to (7), characterized in that the Sm / R ratio inside the main phase grains is lower than the Sm / R ratio in the outer shell portion of the main phase grains. (9) The anisotropic rare earth sintered magnet according to (7) or (8), characterized in that the inside of the main phase grains does not contain Sm. (10) The anisotropic rare earth sintered magnet according to any one of (1) to (9), characterized in that it exhibits a coercive force of 5 kOe or more at room temperature and the temperature coefficient β of the coercive force is -0.5% / K or more. (11) ThMn 12 A method for manufacturing an anisotropic rare earth sintered magnet according to any one of (1) to (10), characterized by pulverizing an alloy containing a compound phase of a ThMn (12) ThMn 12 type crystal, pressure-powder molding in the presence of a magnetic field to form a molded body, and then sintering at a temperature of 800°C or higher and 1400°C or lower. A method for manufacturing an anisotropic rare earth sintered magnet according to any one of (11), characterized by pulverizing and mixing an alloy containing a compound phase of a type crystal and an alloy having a higher R composition ratio and Sm / R ratio than that, and pressure-powder molding in the presence of a magnetic field to form a molded body. (13) ThMn 12 A method for manufacturing an anisotropic rare earth sintered magnet according to (11) or (12), characterized by bringing a material containing Sm into contact with a sintered body having a compound phase of a ThMn type crystal as a main phase and performing heat treatment at a temperature of 600°C or higher and lower than the sintering temperature to diffuse Sm into the inside of the sintered body. (14) The method for manufacturing an anisotropic rare earth sintered magnet according to (13), characterized in that the material containing Sm brought into contact with the sintered body is one or more selected from Sm metal, Sm-containing alloy, Sm-containing compound, and Sm-containing vapor, and the form thereof is one or more selected from powder, thin film, thin strip, foil, and gas. (15) The method for manufacturing an anisotropic rare earth sintered magnet according to any one of (11) to (14), characterized by performing heat treatment on the sintered body at a temperature of 300 to 900°C.

Advantages of the Invention

[0015] According to the present invention, ThMn 12 In an anisotropic rare earth sintered magnet having a ThMn

Brief Description of the Drawings

[0016]

Figure 1

Figure 2

Figure 3

Modes for Carrying Out the Invention

[0017] Hereinafter, embodiments of the present invention will be described. The anisotropic rare earth sintered magnet of the present invention has a composition represented by the following formula (R 1-a Zr a ) x (Fe 1-b Co b ) 100-x-y (M 1 1-c M 2 c ) y and is an anisotropic sintered magnet in which a ThMn 12 type crystal compound is the main phase, contains 80% by volume or more of the main phase composed of the ThMn 12 type crystal compound, the average crystal grain size of the main phase is 1 μm or more, and a two-particle intergranular phase is formed between adjacent main phase grains. Here, x, y, a, b, and c are 7 ≤ x ≤ 15 atomic %, 4 ≤ y ≤ 20 atomic %, 0 ≤ a ≤ 0.2, 0 ≤ b ≤ 0.5, and 0 ≤ c ≤ 0.9, respectively. Thus, since the composition range is wide, it is easy to produce the anisotropic rare earth sintered magnet of the present invention with good reproducibility. First, each component will be described below.

[0018] R is one or more elements selected from rare earth elements, and Sm is essential. Specifically, R may be a combination of Sm and one or more elements selected from Sc, Y, La, Ce, Pr, Nd, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu, with Sm being essential. R is an element necessary for forming a compound having a ThMn 12 type crystal structure. The content of R is 7 atomic % or more and 15 atomic % or less. It is more preferable if it is 8 atomic % or more and 12 atomic % or less. If it is less than 7 atomic %, the α-Fe phase precipitates and sintering becomes difficult. On the other hand, if it exceeds 15 atomic %, the volume ratio of the ThMn 12 type compound phase decreases and good magnetic properties cannot be obtained. The ThMn 12 type compound exhibits a particularly high anisotropy magnetic field H A when R is Sm. Therefore, Sm is essential for the anisotropic rare earth sintered magnet of the present invention. When there is no difference in the Sm concentration between the inside and the outer shell of the main phase grains, the Sm contained in R is preferably 5% or more, more preferably 10% or more, and particularly preferably 20% or more in terms of atomic ratio of R. When the Sm ratio is within such a range, the increasing effect of H A becomes sufficient and a high coercive force can be obtained.

[0019] On the other hand, since the production amount of Sm is less than that of Y, La, Ce, Pr, Nd, etc. and there are resource constraints, it is preferable to utilize Sm as effectively as possible. Therefore, as a tissue form in which Sm is concentrated in the outer shell of the main phase grains, a high coercive force may be obtained with a smaller Sm content. When there is a structure in which the Sm concentration is different between the inside and the outer shell of the main phase grains in this way, the Sm contained in R is preferably 0.1 atomic % or more and 50 atomic % or less in terms of atomic ratio of R. It is more preferably 0.2 atomic % or more and 40 atomic % or less, and particularly preferably 0.5 atomic % or more and 30 atomic % or less. It is more preferable if R is a combination of Sm and one or more elements selected from Y, La, Ce, Pr, Nd.

[0020] Zr brings an effect of enhancing the phase stability by substituting R in the ThMn 12 type compound. The Zr substituting R is 20% or less in terms of atomic ratio of R. If it exceeds 20%, the ThMn 12H of the type compound A decreases, and it is difficult to obtain a high coercive force.

[0021] ThMn 12 It is known that in order for the type crystal structure to stably exist, a third element M is required together with R and Fe as constituent elements. In the anisotropic rare earth sintered magnet of the present invention, M 1 is one or more elements selected from the group consisting of V, Cr, Mn, Ni, Cu, Zn, Ga, Al, and Si, and plays the role of this third element. M 1 is an element that tends to form a compound with R more easily than Fe or has a tendency to be less likely to bond with either Fe or R, compared to M 2 described later that also acts as a third element. One of the features of the anisotropic rare earth sintered magnet of the present invention is that in the magnet structure, together with the ThMn 12 type compound, which is the main phase, an R-rich phase and an R(Fe,Co) 2 phase exist at the grain boundary portion. However, by selecting the element M 1 as the third element, a structure in which these three phases stably coexist can be easily obtained. M 1 and M 2 are collectively denoted as M. M 1 shall occupy at least 10% or more of M in atomic ratio. More preferably, it is 30% or more, and even more preferably, it is 50% or more. If M 1 is less than 10%, the R-rich phase among the above three phases is not stably formed. Also, M 1 and M 2 in total, M, shall be 4 atomic % or more and 20 atomic % or less. If M is less than 4 atomic %, the main phase of the ThMn 12 type compound is not sufficiently formed, and if it exceeds 20 atomic %, the formation amount of the heterogeneous phase increases and good magnet characteristics are not exhibited. The R-rich phase is a phase with a higher concentration of rare earth elements than the main phase. Also, the R(Fe,Co) 2 phase has a MgCu 2 structure and is a compound phase called a Laves phase.

[0022] M 2is one or more elements selected from Ti, Nb, Mo, Hf, Ta, and W. M 2 also ThMn 12 has the effect of stabilizing the type crystal structure, but if it is contained in excess, M 2 carbides such as C phase and MgZn 2 type compound (Fe, Co) 2 M 2 phase precipitates in the main phase or at the grain boundaries. In particular, (Fe, Co) 2 M 2 phase may have a composition richer in Fe than the stoichiometric composition and exhibit ferromagnetic properties like, for example, the Fe 2 Ti phase, which has an adverse effect on the magnetic properties of the sintered magnet. Also, when only M 1 is included without including the third element M 2 alone, it is difficult to stably form the R-rich phase. Therefore, in the case of a composition containing M 2 , its content should be at most 90% by atomic ratio of M.

[0023] The anisotropic rare-earth sintered magnet of the present invention has Sm, M 1 and Fe as essential constituent elements. Furthermore, a part of Fe may be substituted with Co. The substitution with Co increases the Curie temperature T 12 of the ThMn c type compound, which is the main phase, and has the effect of increasing the saturation magnetization M s . The substitution rate of Co should be 50% or less by atomic ratio. If the substitution rate exceeds 50%, M s will conversely decrease. The ratio of Fe and Co is the remainder of R, Zr, M 1 and M 2 . However, in addition to this, inevitable impurities incorporated from raw materials or mixed in the manufacturing process, specifically H, B, C, N, O, F, P, S, Mg, Cl, Ca, etc., may be contained in a total amount of up to 3% by weight.

[0024] 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 R(Fe, Co, M) having a ThMn 12 type crystal structure 12It consists of a compound. Elements such as C, N, and O that are inevitably mixed in the process of manufacturing a sintered magnet are preferably not contained in the main phase. However, in the compositional analysis using EPMA (electron probe microanalyzer), when C, N, and O elements are detected due to measurement variations, the adjustment method of the observation sample, the influence of detection signals of other elements, etc., the upper limit of H in the main phase A is preferably up to 1 atomic% each from the viewpoint of obtaining good properties. The average crystal grain size of the main phase is 1 μm or more, preferably 1 μm or more and 30 μm or less. A range of 1.5 μm or more and 20 μm or less is more preferable, and a range of 2 μm or more and 10 μm or less is particularly preferable. By setting the average crystal grain size within such a range, the decrease in the remanent magnetic flux density B r due to the decrease in the degree of orientation of crystal grains and the decrease in the coercivity H cJ can be suppressed. The volume ratio of the main phase is 80% by volume or more with respect to the entire magnet, preferably 80% by volume or more and less than 99% by volume, and more preferably 90% by volume or more and 95% by volume or less from the viewpoint of obtaining good B r and H cJ . The average crystal grain size of the main phase is the value measured as follows. After polishing the cross-section of the sintered magnet until it becomes mirror-like, it is immersed in an etching solution (such as a mixed solution of nitric acid + hydrochloric acid + glycerin) to selectively remove the grain boundary phase, and observation is performed 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 taken as the average crystal grain size. The volume ratio of the main phase is the value measured as follows. Using EPMA, microstructure observation and compositional analysis of each phase of the anisotropic rare earth sintered magnet are performed to confirm the main phase, R-rich phase, and R(Fe,Co) 2 phase. Then, the volume ratio of each phase is calculated as being equal to the area ratio in the image of the backscattered electron image.

[0025] In order to effectively utilize Sm, a structure may be adopted in which Sm is concentrated in the outer shell portion of the main phase grains, and there are grains in which the Sm concentration inside the main phase grains is lower than that. In that case, although the thickness of the high-Sm outer shell portion is not particularly limited, from the viewpoint of sufficiently obtaining the effect of suppressing the nucleation of reverse magnetic domains in the outer shell portion of the main phase grains and suppressing the reduction in the reduction effect of Sm due to an increase in the Sm content of the entire sintered body, 1 nm to 2 μm is preferable, and 2 nm to 1 μm is particularly preferable. Such a form is caused by increasing the Sm / R ratio (atomic ratio of Sm to R) in the R-rich phase and the R(Fe,Co) 2 phase higher than the Sm / R ratio inside the main phase grains. A structure in which the inside of the main phase grains does not contain Sm is more preferable. Also, the main phase grains having a uniform Sm concentration distribution may be partially included.

[0026] The R-rich phase and R(Fe,Co) 2 phase are formed at the grain boundaries of the magnet structure. The grain boundaries include grain boundary triple points in addition to the inter-particle grain boundary phase. Here, the R-rich phase is a phase containing 40 atomic% or more of R. The present inventors have found that when the above composition containing the M 1 element is used, a magnet containing three phases of the main phase, the R(Fe,Co) 2 phase, and the R-rich phase is easily obtained. For example, in a sintered magnet of the Sm-Fe-Ti ternary system that does not contain the M 1 element, there is a composition region where the Sm(Fe,Ti) 12 main phase and three phases of SmFe 2 , Fe 2 Ti (excluding oxides, etc.) are in equilibrium, but since the Sm(Fe,Ti) 12 main phase and the Sm-rich phase are difficult to be in equilibrium at a low temperature of 400 °C or lower, the Sm-rich phase is not formed as a stable phase. On the other hand, in the case of the Sm-Fe-V ternary system using V, which is one of the M 1 elements, an Fe-V binary compound is not formed, and instead, a Sm-rich phase with a high Sm concentration is formed, and a magnet having three phases of Sm(Fe,V) 12 , SmFe 2 and the Sm-rich phase can be obtained. Also, M 1 , M 2In the Sm-Fe-V-Ti quaternary system containing both, Sm(Fe,V,Ti) 12 , Fe 2 (V,Ti), SmFe 2 and four phases of the Sm-rich phase can stably exist. In the anisotropic rare-earth sintered magnet of the present invention, based on such findings, a composition containing a predetermined amount of M 2 element is selected to form an R-rich phase and an R(Fe,Co) 1 phase at the grain boundary portion.

[0027] 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 to form a liquid phase, so liquid-phase sintering proceeds, and sintering is completed more quickly 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 quickly.

[0028] 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 ThMn 12 type compound crystal grains during the sintering process or the subsequent aging process, and this cleaning effect suppresses the nucleation of reverse magnetic domains, which is a factor in reducing coercivity. In particular, the R(Fe,Co) 2 phase has relatively high wettability with respect to the ThMn 3 , RM 2 , R(Fe,Co)M, R(Fe,Co) 2 M 2 and other compound phases compared to the ThMn 12 phase, and is likely to coat the surface of the main-phase grains, so the cleaning effect is large.

[0029] The third effect is the formation of an intergranular phase between two particles. In a magnet containing an R-rich phase in the structure, by performing an appropriate sintering process or aging treatment, an intergranular phase containing more R than the main phase is formed between adjacent ThMn 12 type compound main phase grains. As a result, the magnetic interaction between the main phase grains is weakened, and the sintered magnet exhibits a high coercive force. However, since the composition region in which only the two phases of the ThMn 12 type compound main phase and the R-rich phase are in equilibrium is extremely limited, it is difficult to stably manufacture such a magnet considering the composition variation. ThMn 12 type compound main phase, R-rich phase, and R(Fe,Co) 2 phase, a structure in which the surface of the main phase grains is covered by the intergranular phase between two particles can be stably formed. Also, in a magnet without an R-rich phase, it is difficult to form an intergranular phase between two particles or to cover the surface of the main phase grains with the intergranular phase between two particles, so it is difficult to obtain a magnet exhibiting sufficient coercive force.

[0030] The fourth effect is to increase the Sm concentration at the grain boundaries. In order to have a structure with different Sm concentrations in the interior and outer shell of the main phase grains, when applying the grain boundary diffusion method as a manufacturing method, the R-rich phase and R(Fe,Co) 2 phase existing at the grain boundaries become a liquid phase during the diffusion treatment and play a role in diffusing and penetrating Sm placed on the sintered body into the interior. Therefore, the Sm / R ratio in at least one of the R-rich phase or R(Fe,Co) 2 phase becomes higher than the Sm / R ratio inside the main phase grains. Also, when applying the two-alloy method as a manufacturing method, by using an alloy mainly composed of the ThMn 12 type compound phase and an alloy having a higher R composition ratio and Sm / R ratio than that, the Sm / R ratio in at least one of the R-rich phase or R(Fe,Co) 2 phase of the sintered body becomes higher than the Sm / R ratio inside the main phase grains. The concentration of Sm in the R-rich phase and R(Fe,Co) 2 phase increases the Sm concentration in the outer shell of the main phase grains in contact with these grain boundary phases, and H A is improved and the coercive force of the sintered magnet increases.

[0031] The R-rich phase shall contain at least 40 atomic % or more of R as described above. If R is less than 40 atomic %, the wettability with the main phase is insufficient, and thus it is difficult to obtain the above-described effects. It is more preferable that R contains 50 atoms or more, and particularly preferably 60 atoms or more. The R-rich phase may be an R metal phase such as the above-described Sm phase, an amorphous phase, or R 3 (Fe, Co, M), R 2 (Fe, Co, M), R 5 (Fe, Co, M) 3 It may also be an intermetallic compound with a high R composition and a low melting point such as R(Fe, Co, M). Further, impurity elements such as Fe, Co, M elements, and H, B, C, N, O, F, P, S, Mg, Cl, Ca may be contained up to a total of 60 atomic %.

[0032] On the other hand, the R(Fe, Co) 2 phase is a Laves compound of the MgCu 2 type crystal. However, when performing composition analysis using EPMA or the like, considering measurement variations and the like, it shall contain R of 20 atomic % or more and less than 40 atomic %. Also, a part of Fe and Co may be substituted by the M element. However, the substitution amount of M shall be within the range where the MgCu 2 type crystal structure is maintained.

[0033] 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 means a phase that exhibits ferromagnetic or ferrimagnetic properties and has a Curie temperature T c of 23 °C or higher at room temperature. RFe 2 has a T 2 of 23 °C or higher except for CeFe c , and CeFe 2 also has a T c of 23 °C or higher if 10% or more of R is substituted by other elements. On the other hand, RCo 2 has a T 2 of 23 °C or lower or is a paramagnetic phase except for GdCo c . However, in the anisotropic rare earth sintered magnet of the present invention, since the substitution atomic ratio of Fe by Co is 0.5 or less, in most cases, R(Fe, Co) 2The phase becomes a magnetic phase. Generally, the soft magnetic phase contained in the structure often has an adverse effect on magnetic properties. However, in the anisotropic rare earth sintered magnet of the present invention, the cleaning effect of the surface of the main phase grains by the R(Fe,Co) 2 phase and the effect of forming the interparticle grain boundary phase between two particles are greater, and even if it is a magnetic phase, it is considered to contribute to an increase in coercivity.

[0034] The R-rich phase and R(Fe,Co) 2 The total formation amount of the phase is preferably 1% by volume or more, and more preferably 1% by volume or more and less than 20% by volume. Further, 1.5% by volume or more and less than 15% by volume is more preferable, and the range of 2% by volume or more and less than 10% by volume is even more preferable. By setting such a range, the area in contact with the main phase grains is ensured, and the effect of increasing H cJ is easily obtained. Also, the decrease in B r is suppressed, and the desired magnetic properties are easily obtained.

[0035] In the anisotropic rare earth sintered magnet of the present invention, as described above, an R-rich phase and an R(Fe,Co) 2 phase exist at the grain boundary portion, and an interparticle grain boundary phase is formed between adjacent main phase grains composed of a ThMn 12 type compound. By covering the surface of the main phase grains with the interparticle grain boundary phase, the magnetic interaction between the main phase grains is weakened, and a high coercivity is exhibited.

[0036] The above interparticle grain boundary phase may be in an amorphous state with a disordered atomic arrangement or may have regularity in the atomic arrangement. Also, the R-rich phase and R(Fe,Co) 2It may be the same phase. When observing the grain boundary phase between two particles using an apparatus such as STEM (scanning transmission electron microscope), its composition preferably contains 20 atomic% or more of R. By setting it within such a range, the magnetic coupling between the main phase grains can be sufficiently reduced, and a high coercive force is easily obtained. Also, the thickness of the grain boundary phase between two particles is preferably 0.5 nm or more. Thereby, it is easy to ensure the magnetic separation effect between the main phase grains, and a sufficient coercive force improvement effect can be obtained. Furthermore, its thickness is preferably 1 μm or less, more preferably 0.5 μm or less, and even more preferably 100 nm or less. If it is within such a range, it is easy to suppress the influence of the deterioration of magnetic properties due to the decrease in the volume ratio of the main phase grains from becoming larger than the effect of increasing the coercive force. Incidentally, the thickness of the grain boundary phase between two particles was measured from the STEM image as follows. Using a STEM apparatus (JEM-ARM200F manufactured by JEOL Ltd.), at least three or more locations of one sample were observed at locations where adjacent main phase grains were in contact. The thickness of the grain boundary phase between two particles was measured from the HAADF (High-Angle Annular Dark Field) image obtained by the observation, and the average value of these thicknesses was taken as the thickness of the grain boundary phase between two particles.

[0037] In addition, the anisotropic rare earth sintered magnet of the present invention may contain R oxides, R carbides, R nitrides, M carbides, etc. formed by C, N, O inevitably mixed in. From the viewpoint of suppressing the deterioration of magnetic properties, the volume ratio of these is preferably 10 volume% or less, more preferably 5 volume% or less, and particularly preferably 3 volume% or less.

[0038] It is preferable that the phases other than the above are as few as possible. For example, R 2 (Fe, Co, M) 17 phase, R 3 (Fe, Co, M) 29 When the phase exists in the magnet structure, from the viewpoint of suppressing the influence on magnetic properties and the resulting decrease in coercive force, the formation amount thereof is preferably less than 1 volume% each. Also, from the viewpoint of ensuring a sufficient proportion of the main phase, (Fe, Co) 2 M phase or RM where R is less than 40 atomic% 3 , RM2 , R(Fe, Co)M, R(Fe, Co) 2 M 2 etc. are each preferably less than 1% by volume. The total of these phases is preferably 3% by volume or less. Further, from the viewpoint of preventing a significant decrease in magnetic properties, the α-(Fe, Co) phase is preferably not contained in the anisotropic rare earth sintered magnet of the present invention.

[0039] Next, the manufacturing method will be described. The anisotropic rare earth sintered magnet of the present invention is manufactured by powder metallurgy. First, in order to produce a raw material alloy, metal raw materials, alloys, ferroalloys, etc. of R, Fe, Co, and M are used, and after considering raw material losses during the manufacturing process, etc., it is adjusted so that the finally obtained sintered body has 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 produce the alloy by adjusting the cooling rate 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 do not sufficiently orient in the magnetic field forming process, leading to a decrease in B r . When α-Fe precipitates in the alloy, heat treatment may be performed on the alloy so that the amount of formation of the ThMn 12 type compound phase increases. The alloy may be a single composition alloy, or may be adjusted by a method of preparing a plurality of alloys with different compositions and mixing their powders in a later process.

[0040] The above raw material alloy is roughly pulverized into a powder with an average particle size of 0.05 to 3 mm by means such as mechanical pulverization using a brown mill or hydrogenation pulverization. Alternatively, the HDDR method (hydrogen disproportionation desorption recombination method) used as a method for manufacturing Nd-Fe-B magnets may be applied. Further, the coarse powder is finely pulverized by a ball mill or a jet mill using high-pressure nitrogen or the like to obtain a powder with an average particle size of 0.5 to 20 μm, more preferably 1 to 10 μm. Note that lubricants or the like may be added before and after the fine pulverization process as necessary. Next, using a magnetic field pressing device, the powder is formed while orienting the easy magnetization axis of the alloy powder in an applied magnetic field to obtain a compacted body. The forming is preferably performed in a vacuum, a nitrogen gas atmosphere, an inert gas atmosphere such as Ar, etc. to suppress oxidation of the alloy powder.

[0041] The step of sintering the compacted body is carried out at a temperature of 800°C or higher and 1400°C or lower in a vacuum or an inert atmosphere using a sintering furnace. If it is less than 800°C, sintering does not proceed sufficiently and a high sintering density cannot be obtained. If it exceeds 1400°C, the main phase of the ThMn 12 type compound decomposes and α-Fe precipitates. The sintering temperature is particularly preferably in the range of 900 to 1300°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 may be used in which, after heating to a first sintering temperature, it is held at a lower second sintering temperature for a predetermined time to refine the crystal grains. Also, sintering may be performed multiple times, or a method such as spark plasma sintering 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 coercivity, an aging heat treatment may be further performed at 300 to 900°C for 0.5 to 50 hours. By optimizing the sintering and aging conditions according to the composition, powder particle size, etc., an improvement in H cJ is brought about. Further, the sintered body is cut and ground into a predetermined shape and magnetized to obtain a sintered magnet.

[0042] On one hand, there are anisotropic rare-earth sintered magnets in which there are main-phase grains with an Sm / R ratio inside the main-phase grains that is lower than the Sm / R ratios of the R-rich phase and the R(Fe,Co) 2 Examples of means for manufacturing such magnets include, for example, the two-alloy method and the grain boundary diffusion method.

[0043] When using the two-alloy method, metal raw materials, alloys, ferroalloys, etc. of R, Fe, Co, and M are used to produce two raw material alloys with different compositions. Note that three or more alloys may also be used. At this time, ThMn 12 It is preferable to combine alloy A mainly composed of the ThMn-type compound phase with a relatively low Sm / R ratio and alloy B with a relatively higher R composition ratio and Sm / R ratio than alloy A, and adjust so that the average composition becomes a predetermined composition. These alloys are produced by a casting method or a strip casting method and pulverized. The step of mixing the respective alloy powders may be performed in the coarse powder state before fine pulverization or after fine pulverization. Further, molding and sintering are performed to obtain a sintered body. Aging heat treatment may be performed to improve the coercive force.

[0044] In the sintered magnet by the two-alloy method, a main phase composed mainly of the components of alloy A forms a ThMn 12 type compound, and an R-rich phase, an R(Fe,Co) 2 phase and the outer shell part of the main-phase grains are mainly formed by the components of alloy B. Therefore, the Sm / R atomic ratio of the R-rich phase and the R(Fe,Co) 2 phase formed at the grain boundary part becomes higher than the Sm / R atomic ratio inside the main-phase grains. Also, a part of Sm in the grain boundary phase substitutes R atoms in the surface layer part of the main-phase grains, forming a core-shell structure with different Sm concentrations between the grain surface layer part and the inside, and increasing the coercive force.

[0045] On the other hand, in the grain boundary diffusion method, first, a sintered body is produced in the same manner as described above by the single-alloy method or the two-alloy method. At this time, R in the composition of the sintered body may contain Sm or may not contain Sm.

[0046] Next, grain boundary diffusion of Sm is performed on the obtained sintered body. After cutting and grinding the sintered body as needed, a diffusion material selected from compounds such as metals, alloys, oxides, fluorides, oxyfluorides, hydrides, carbides, etc. containing Sm 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 preferably, particularly 20 to 500 μg / mm 2 is preferred. Within such a range, an increase in H cJ can be sufficiently obtained, and an increase in manufacturing cost due to an increase in the Sm content can be suppressed. Also, taking advantage of the property that Sm has a high vapor pressure, Sm metal or Sm alloy may be heat-treated together with the sintered body in the same chamber and brought into contact with the sintered body as Sm vapor.

[0047] This sintered body is heat-treated in a vacuum or an inert gas atmosphere with Sm placed on its surface. The heat treatment temperature is preferably 600°C or higher and the sintering temperature or lower, particularly preferably 700°C or higher and 1100°C or lower. The heat treatment time is preferably 0.5 to 50 hours, 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. In order to improve the coercivity, an aging heat treatment may be further performed at 300 to 900°C for 0.5 to 50 hours.

[0048] Sm arranged on the sintered body penetrates into the sintered body while increasing the Sm concentration in the R-rich phase and the R(Fe,Co) 2 phase, and the Sm / R ratio of these grain boundary phases increases. As the Sm concentration of the grain boundary phase increases, substitution of R atoms by Sm also occurs in the surface layer portion of the main phase grains in contact with the grain boundary phase, and the Sm / R ratio of the surface layer portion of the main phase grains becomes higher than the Sm / R ratio inside the main phase grains, resulting in an increase in H cJ .

[0049] The anisotropic rare earth sintered magnet of the present invention thus produced has a residual magnetic flux density B of 5 kG or more at room temperaturer and a coercive force H of at least 5 kOe or more cJ is shown. Room temperature H cJ is more preferably 8 kOe or more. Further, the temperature coefficient β of the coercive force exhibits a characteristic of -0.5% / K or more. Here, β = ΔH cJ / ΔT × 100 / H cJ (20 °C) (ΔH cJ = H cJ (20 °C) - H cJ (140 °C), ΔT = 20 - 140 (°C)). The anisotropic rare earth sintered magnet of the present invention has a smaller temperature change in coercive force compared to the Nd-Fe-B sintered magnet and is suitable for use at high temperatures. [Example]

[0050] 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.

[0051] [Example 1] The composition was adjusted using Sm metal, electrolytic iron, ferrovanadium, Al metal, and Si, and melted in an Ar gas atmosphere by a high-frequency induction furnace to produce a cast alloy. In order to disappear the primary crystal α-Fe, the alloy was heat-treated at 900 °C for 50 hours. The structure of the alloy obtained by a laser microscope (manufactured by Olympus Corporation, LEXT OLS4000) was observed, and it was confirmed from the observed image that the average crystal grain size of the main phase was 5 μm or more. After subjecting the alloy to a hydrogen storage treatment and a dehydrogenation treatment of heating at 400 °C in a vacuum to obtain a coarse powder, it was pulverized by a jet mill in a nitrogen stream to produce a fine powder with an average particle size of 1.8 μm. Further, the fine powder was filled into a 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. After sintering this compacted powder compact at 1140 °C for 3 hours in an Ar gas atmosphere, it was cooled to room temperature at a cooling rate of 13 °C / min to obtain a sintered body.

[0052] The composition of the sintered body analyzed by inductively coupled plasma optical emission spectrometry (ICP-OES) using a high-frequency inductively coupled plasma optical emission spectrometer (manufactured by Hitachi High-Tech Science Corporation, SPS3520UV-DD) was Sm 10.2 Fe bal. V 14.9 Al 0.5 Si 0.2 It was also confirmed by X-ray diffraction that the ThMn 12 type crystal was the main phase. Using an EPMA device (manufactured by JEOL Ltd., JXA-8500F), the microstructure of the sintered body was observed and the composition analysis of the formed phases was performed. It was confirmed that an R-rich phase and an R(Fe,Co) 2 phase were formed at the grain boundary triple points. The volume ratios of the main phase, R-rich phase, and R(Fe,Co) 2 phase were calculated as being equal to the area ratios in the image of the backscattered electron image. As a result, in the grain boundary portion of the sintered body structure, the R-rich phase and the R(Fe,Co) 2 phase each existed at 1% by volume or more. Also, the average crystal grain size of the main phase calculated from the results of observing the sintered body sample after etching was 9.9 μm. Furthermore, the room temperature H cJ measured with a B-H tracer was 8.5 kOe, and the temperature coefficient β of H cJ was -0.46% / K.

[0053] From this sintered body, an observation sample was thinly cut out using a FIB-SEM device (Scios Dual Beam manufactured by FEI) and observed with a STEM device (JEM-ARM200F manufactured by JEOL Ltd.). The obtained HAADF (High-Angle Annular Dark Field) images are shown in FIGS. 1 and 2. From FIG. 1, it can be confirmed that an intergranular phase 13 exists at the grain boundary portion sandwiched between two main phase grains 11 and 12. The thickness of the intergranular phase 13 between the two particles at this time was about 0.7 nm, and the composition obtained by EDX analysis (energy-dispersive X-ray fluorescence analysis) was Sm 61.2 Fe bal. V 7.1 Al 1.3 Si 1.1 It was. Also, in FIG. 2 where another location of the same sintered body was observed, the intergranular phase 13 between the two particles had a thickness of about 1.4 nm, and the composition was Sm31.9 Fe bal. V 9.5 Al 2.7 Si 0.6 It was. The results are shown in Tables 1 to 3 and 5.

[0054] [Comparative Example 1] Using Sm metal, electrolytic iron, and Ti metal, the composition was adjusted. In the same manner as in Example 1, a cast alloy was produced by a high-frequency induction furnace, and further heat-treated at 900 °C for 50 hours. The microstructure of the alloy obtained by a laser microscope was observed, and it was confirmed from the observed image that the average crystal grain size of the main phase was 5 μm or more. In the same manner as in Example 1, pulverization and molding in a magnetic field were performed, and after sintering at 1175 °C for 3 hours in an Ar gas atmosphere, it was cooled to room temperature at a cooling rate of 13 °C / min to obtain a sintered body of Comparative Example 1. The composition values of the sintered body analyzed by the ICP method were Sm 9.7 Fe bal. Ti 8.1 It was. Also, from X-ray diffraction measurement, it was confirmed that the main phase of Comparative Example 1 was a ThMn 12 type crystal. When the formed phases were examined by EPMA, the R(Fe,Co) 2 phase was present, but the R-rich phase was not formed, and fine TiC phases were precipitated. When the magnetic properties were measured with a B-H tracer, this Comparative Example 1 showed only a low coercive force of 0.1 kOe at room temperature. The HAADF image obtained in Comparative Example 1 is shown in Fig. 3. At the boundary between the two main phase grains 11 and 12, the inter-particle grain boundary phase as seen in Example 1 was not formed. The results are shown in Tables 1 to 3.

[0055] [Examples 2 to 8] Similar to Example 1, the composition was adjusted and a cast alloy was produced by high-frequency melting. To eliminate primary α-Fe, the alloy was heat-treated at 850 - 1100 °C for 10 - 50 hours. The microstructure of the alloy obtained by a laser microscope was observed, and it was confirmed that the average crystal grain size of the main phase was 1 μm or more in all the observed images. After performing hydrogen storage treatment and dehydrogenation treatment by heating at 450 °C in vacuum to obtain coarse powder, it was pulverized with a jet mill in a nitrogen stream to produce fine powder with an average particle size of 2 - 4 μm. Further, the fine powder was filled into the mold of a molding device in an inert gas atmosphere and molded in a magnetic field. After sintering this compacted body in an Ar gas atmosphere, it was cooled to room temperature, and further aging heat treatment was performed to obtain a sintered body sample. Table 1 shows the composition of each sample analyzed by the ICP method, the crystal structure of the main phase confirmed by X-ray diffraction, and the average crystal grain size of the main phase of the sintered body. Table 2 shows the sintering treatment conditions, the cooling rate after sintering, the aging treatment conditions, and the B r , H cJ , and H cJ temperature coefficient β of each example. In Example 8, a two-step sintering method was applied in which the temperature was lowered to the second sintering temperature immediately after rising to the first sintering temperature and held for a predetermined time. Table 3 also shows the composition and phase ratio of each phase analyzed by EPMA. In the samples of Examples 2 - 8, an R-rich phase and an R(Fe,Co) 2 phase were formed in the microstructure, showing a coercive force of 5 kOe or more at room temperature and a temperature coefficient β of -0.5% / K or more. Further, when STEM observation was performed on these sintered body samples in the same manner as in Example 1, it was confirmed that an intergranular phase between two particles existed at the grain boundary portion sandwiched between two main phase grains in all examples. Table 5 shows the composition and thickness of the measured intergranular phase between two particles.

[0056] [Comparative Examples 2 - 6] Sintered body samples of Comparative Examples 2 to 5 were prepared in the same manner as in Example 2, except that the composition was adjusted to that shown in Table 1. The results are shown in Tables 1, 2, and 4. In Comparative Example 2, the total amount of R was less than 7 atomic %, and sufficient sintering could not be achieved. A large amount of α-Fe phase was formed in the sintered body. In Comparative Example 3, the total amount of R exceeded 15 atomic %, and the volume ratio of the main phase was less than 80%. In Comparative Example 4, the total amount of M element exceeded 20 atomic %, no R-rich phase was observed, and the RFeSi phase of the PbClF-type crystal was formed. In Comparative Example 5, the RCu 2 phase of the KHg 2 type crystal was present at the grain boundary triple point, but the total amount of M element exceeded 20 atomic %, and no R-rich phase was found. In Comparative Example 6, the total amount of M was less than 4 atomic %, no ThMn 12 type crystal was observed in the structure, and the main phase of the Th 2 Zn 17 type crystal was formed.

[0057] [Comparative Example 7] The composition was adjusted using Sm metal, electrolytic iron, Ti metal, and V metal, and the molten raw material was cooled on a rotating Cu roll at a peripheral speed of 20 m / sec to produce a rapidly quenched ribbon raw material alloy. The thickness of the ribbon was 10 to 50 μm, and the structure of the alloy obtained by a laser microscope was observed. Although the average crystal grain size was too fine to measure from the observed image, it was confirmed to be less than at least 1 μm. After pulverizing this alloy ribbon with a ball mill, powders of 300 μm or less were selected by a sieve, and hot pressing was performed at 750 °C in an Ar atmosphere. The average crystal grain size of the main phase grains was as fine as about 0.2 to 0.3 μm, and the compositions of the main phase and the grain boundary phase could not be identified by EPMA. Also, since the easy magnetization axes of the main phase were not aligned, only a low B r was obtained. The results are shown in Tables 1, 2, and 4.

[0058] [Example 9] The composition was adjusted using Nd metal, Y metal, electrolytic iron, pure Si, and Hf metal. After melting in an Ar gas atmosphere using a high-frequency induction furnace, strip casting was performed on a water-cooled Cu roll to produce a rapidly solidified ribbon alloy with a thickness of about 0.2 to 0.4 mm and a composition of 7.5 atomic% Nd, 1.0 atomic% Y, 13.0 atomic% Si, 1.0 atomic% Hf, and the balance Fe. The average crystal grain size in the minor axis direction of the alloy determined from the image observed with a laser microscope was 2.5 μm. After performing a hydrogen storage treatment on this alloy at room temperature, a dehydrogenation treatment was carried out by heating in a vacuum at 400 °C to obtain coarse powder (designated as Sample 9A powder). On the other hand, an alloy ingot with a composition of 40 atomic% Sm, 10 atomic% Ga, 5 atomic% Cu, and the balance Co was produced using Sm metal and electrolytic iron as raw materials in a high-frequency induction furnace, and then made into coarse powder by mechanical pulverization (designated as Sample 9B powder). After mixing Sample 9A powder and Sample 9B powder at a weight ratio of 95:5, they were pulverized with a jet mill in a nitrogen stream to produce fine powder with an average particle size of 1.8 μm.

[0059] Using this mixed powder, molding in a magnetic field was performed in the same manner as in Example 1. After sintering at 1200 °C for 3 hours in an Ar gas atmosphere, it was cooled to room temperature at a cooling rate of 12 °C / min, and further heat-treated at 650 °C for 1 hour in an Ar gas atmosphere to obtain the sintered body of Example 9. The composition values of the sintered body sample were Sm 1.8 Nd 7.2 Y 1.0 Fe bal. Co 1.0 Si 12.8 Ga 0.6 Cu 0.4 Hf 1.0 It was. Also, from X-ray diffraction measurement, it was confirmed that the main phase of this sintered body was a ThMn 12 type crystal. The composition of the main phase measured by EPMA was that the central part of the grain was Nd 6.4 Y 1.1 Fe bal. Co 1.0 Si 12.7 Ga 0.5 Cu 0.1 Hf 1.1 and did not contain Sm, but Sm was present in the outer shell part of the grain 3.5 Nd 3.0 Y 1.0 Fe bal. Co1.0 Si 13.0 Ga 0.4 Cu 0.1 Hf 0.9 and it was confirmed that the Sm / R ratio inside the grains was lower than the Sm / R ratio on the surface layer. Further, the microstructure of the sintered body was observed by EPMA and the composition analysis of each phase was carried out, and it was confirmed that an R-rich phase and R(Fe,Co) 2 phase were present in an amount of 1% by volume or more at the grain boundaries. In addition, a slight amount of the RCu 2 phase was confirmed. The R 2 (Fe,Co,M) 17 phase, the R 3 (Fe,Co,M) 29 phase, and the α-Fe phase were not observed. Since there are also phases such as oxides, the total of the phase ratios is less than 100%.

[0060] The R-rich phase, R(Fe,Co) 2 phase, and the RCu 2 phase composition analysis values were Sm 26.8 Nd 31.7 Y 0.1 Fe bal. Si 36.2 Ga 5.2 , Sm 17.2 Nd 17.4 Y 0.2 Fe bal. Co 0.4 Si 0.3 Hf 0.1 , Sm 15.9 Nd 18.6 Fe bal. Cu 65.2 respectively. From this, it was confirmed that the Sm / R ratio inside the grains was lower than the Sm / R ratio of the R-rich phase and the R(Fe,Co) 2 phase. The average crystal grain size of the main phase was 8.6 μm. The coercive force of this sintered body was 5.6 kOe at room temperature, and the temperature coefficient β of the coercive force was -0.45% / K. The Curie temperature T 2 of the alloy with the same composition prepared based on the analysis value of the R(Fe,Co) c phase was 318°C. Regarding Example 9, STEM observation was performed in the same manner as in Example 1, and it was confirmed that an intergranular phase between two particles was present at the grain boundary sandwiched between two main phase grains. The composition of the measured intergranular phase between two particles was Sm21.7 Nd 24.5 Fe bal. Co 0.5 Si 12.8 Ga 2.6 Cu 8.0 and the thickness was 35 nm.

[0061] [Example 10] The composition was adjusted using Ce metal, La metal, electrolytic iron, Co metal, pure Si, and Mo metal. After melting in an Ar gas atmosphere using a high-frequency induction furnace, strip casting was performed on a water-cooled Cu roll to produce an alloy ribbon with a thickness of about 0.2 to 0.4 mm. When the average grain boundary spacing of this alloy was calculated, it was 4.1 μm. The alloy was subjected to the same hydrogen storage treatment and dehydrogenation treatment as in Example 9 to obtain coarse powder, which was further pulverized with a jet mill in a nitrogen stream to produce fine powder with an average particle size of 2.9 μm. Next, the fine powder was pressure-molded while being oriented in a magnetic field, sintered in a vacuum at 950 °C for 1.5 hours, and then cooled to room temperature at a cooling rate of 11 °C / min and taken out to obtain a sintered body. This sintered body was placed in a vacuum heat treatment furnace together with Sm metal and heat-treated at 780 °C for 8 hours. After being taken out of the furnace once, it was further subjected to an aging treatment at 520 °C for 2 hours to obtain Example 10.

[0062] As a result of ICP analysis of the sintered body sample of Example 10, the composition was Sm 2.4 Ce 7.7 La 1.1 Fe bal. Co 0.6 Si 12.6 Mo 0.9 . 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 of the ThMn 12 type. Also, the microstructure of the sintered body and the composition analysis of each phase were performed by EPMA, and it was confirmed that an R-rich phase and an R(Fe,Co) 2 phase were present in an amount of 1% by volume or more at the grain boundary. R 2 (Fe,Co,M) 17 phase, R 3 (Fe,Co,M) 29 phase, and α-Fe phase were not observed. Since phases such as oxides also exist, the total of the phase ratios is less than 100%.

[0063] The composition analysis values by EPMA of the central part and the outer shell part of the main phase grains are Ce for each 7.3 La 0.1 Fe bal. Co 0.5 Si 13.5 Mo 1.0 , Sm 3.2 Ce 4.2 La 0.3 Fe bal. Co 0.5 Si 13.5 Mo 1.0 , and it was confirmed that the Sm / R ratio inside the grains was lower than that of the surface layer. Also, the composition analysis values of the R-rich phase and the R(Fe,Co) 2 phase are Sm for each 31.1 Ce 19.6 La 23.2 Fe bal. Co 2.7 Si 0.6 Mo 0.2 , Sm 15.0 Ce 18.9 La 0.5 Fe bal. Co 0.2 Si 0.7 Mo 0.2 . Sm was not detected inside the main phase grains, while the R-rich phase and the R(Fe,Co) 2 phase existing at the grain boundaries contain Sm, and it was confirmed that the Sm / R ratio is high. R(Fe,Co) 2 Based on the analysis values of the phase, an alloy with the same composition was prepared by arc melting. After homogenization treatment at 800 °C for 20 hours, magnetization-temperature measurement was performed by VSM. As a result, the Curie temperature T c was 140 °C. Also, the average crystal grain size of the main phase calculated from the results of observing the sintered body of Example 9 after etching was 12.3 μm. Furthermore, when the magnetic properties were measured with a B-H tracer, the room temperature coercive force H cJ showed 6.3 kOe. Also, the temperature coefficient β of H cJ was -0.48% / K.

[0064] STEM observation was carried out for Example 10 in the same manner as in Example 1, and it was confirmed that there was an intergranular phase between two particles at the grain boundary portion sandwiched between the two main phase grains. The measured composition of the intergranular phase between two particles was Sm 21.2 Ce 15.5 La 25.9 Fe bal. Co 1.0 Si 0.6 and the thickness was 92 nm.

[0065] [Comparative Example 8] A sintered body of Comparative Example 9 was produced by the same production method as the sintered body of Example 10, except that the step of heat-treating with Sm metal was not performed and aging treatment was carried out at 520 °C for 2 hours.

[0066] The composition of the sintered body of Comparative Example 8 did not contain Sm, but was Ce 8.3 La 1.3 Fe bal. Co 0.6 Si 13.0 Mo 0.9 and the composition analysis value of the main phase grains was Ce 7.5 La 0.3 Fe bal. Co 0.6 Si 13.1 Mo 0.9 There was no R(Fe,Co) 2 phase at the grain boundary portion, and two types of R-rich phases with compositions of Ce 33.1 La 29.6 Fe bal. Si 37.3 and Ce 23.3 La 54.3 Fe bal. Co 0.8 Si 0.6 Mo 0.1 were observed. The room temperature coercive force H cJ of Comparative Example 8 was 0.1 kOe. The results are shown in Tables 6 to 9.

[0067]

Table 1

[0068]

Table 2

[0069]

Table 3

[0070]

Table 4

[0071]

Table 5

[0072]

Table 6

[0073]

Table 7

[0074]

Table 8

[0075]

Table 9

Explanation of Symbols

[0076] 11, 12 Main Phase Grains 13 Grain Boundary Phase between Two Particles

Claims

Claim 1 ThMn １２ A method for manufacturing an anisotropic rare-earth sintered magnet, characterized in that an alloy containing a compound phase of a １２ -type crystal, an alloy having a higher R composition ratio and Sm / R ratio than that alloy are pulverized and mixed, compacted under a magnetic field application to form a compact, and then sintered at a temperature of 800°C or higher and 1400°C or lower to obtain a sintered body. Claim 2 An alloy containing a compound phase of ThMn12 type crystals is pulverized, compacted under a magnetic field to form a compact, and then sintered at a temperature of 800 °C or higher and 1400 °C or lower to obtain ThMn １２ A method for manufacturing an anisotropic rare earth sintered magnet, characterized in that a material containing Sm is brought into contact with a sintered body having a compound phase of ThMn type crystals as a main phase, heat treatment is performed at a temperature of 600 °C or higher and lower than the sintering temperature to diffuse Sm into the sintered body, and the Sm / R ratio of the material containing Sm is higher than the Sm / R ratio of the sintered body. Claim 3 The material containing Sm to be brought into contact with the sintered body is one or more selected from Sm metal, Sm-containing alloys, compounds containing Sm, and vapors containing Sm, and its form is one or more selected from powder, thin film, ribbon, foil, and gas. The method for manufacturing an anisotropic rare earth sintered magnet according to claim 2, characterized in that it is one or more. Claim 4 The method for manufacturing an anisotropic rare earth sintered magnet according to any one of claims 1 to 3, characterized in that the sintered body is heat-treated at a temperature of 300 to 900 °C.

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