Anisotropic rare earth sintered magnet and its manufacturing method

Abstract

This anisotropic rare earth sintered magnet has a composition represented by the formula (R1-aZra)x(Fe1-bCob)100-x-y(M1 1-cM2 c)y (therein, R is one or more elements selected from among the rare earth elements and must include Sm, M1 is one or more elements selected from the group consisting of V, Cr, Mn, Ni, Cu, Zn, Ga, Al and Si, M2 is one or more elements selected from the group consisting of Ti, Nb, Mo, Hf, Ta and W, and x, y, a, b and c are such that 7 ≤ x ≤ 15 atom%, 4 ≤ y ≤ 20 atom%, 0 ≤ a ≤ 0.2, 0 ≤ b ≤ 0.5, and 0 ≤ c ≤ 0.9), and is characterized by containing 80 vol% or more of a main phase comprising a ThMn12 type crystal compound, with the average crystal grain diameter of the main phase being 1 µm or more, and an interparticle grain boundary phase being formed between adjacent main phase grains. This method for producing an anisotropic rare earth sintered magnet is characterized by pulverizing an alloy containing a ThMn12 type crystal compound phase, carrying out powder compression molding in a magnetic field so as to obtain a molded body, and then sintering the molded body at a temperature of 800-1400ºC. The present invention makes it possible to provide an anisotropic rare earth sintered magnet which contains a ThMn12 type crystal compound as the main phase and which exhibits good magnetic characteristics; and a method for producing the same.

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 12A rare earth permanent magnet has been reported in which the RT compound has a cubic crystal structure (R is one or more rare earth elements with La as an essential element, and T is Fe, or Fe and Co, or an element in which a part of the Fe and Co elements is replaced with M (one or more elements selected from Ti, V, Cr, Mo, W, Zr, Hf, Nb, Ta, Al, Si, Cu, Zn, Ga, and Ge)), the grain boundary phase has a cubic crystal structure, and the La-rich phase σ with a La composition ratio of 20 at% or more accounts for 20% or more of the cross-sectional area. It is said that the inclusion of a non-magnetic cubic La-rich phase in the grain boundary portion provides magnetic separation between the main phases and reduces the interface strain between the grain boundary phase and the main phase.

[0005] In Patent Document 3, ThMn 12 The main phase has a crystal structure of the Sm type. 5 Fe 17 System phase, SmCo 5 System phase, Sm 2 O 3 Phase system and Sm 7 Cu 3 There have been reports of rare earth magnets with subphases containing any of the Sm phases, with the subphase volume fraction ranging from 2.3 to 9.5%. 5 Fe 17 Phase and SmCo 5 The Sm-based phase is a magnetic phase that exhibits higher magnetic anisotropy than the main phase, and it isolates the crystal grains of the main phase and prevents the movement of domain walls within the main phase, thereby improving the magnetization and coercivity of the magnet. 2 O 3 Phase and Sm 7 Cu 3 The Sm phase is a non-magnetic phase, and by isolating each crystal grain of the main phase, it is believed that the magnetization reversal of the main phase is prevented from propagating to the surroundings, thereby improving the magnetization and coercive force of the magnet. 7 Cu 3 It is stated that the system is a non-equilibrium phase.

[0006] In Patent Document 4, an alloy having a main phase and one or more subphases and a composition of the entire alloy is R(Fe,Co). w-z Ti z Cu α(R is at least one rare earth element, 8≦w≦13, 0.42≦z<0.70, 0.40≦α≦0.70) have been reported. In addition, the subphase is mainly composed of a crystalline phase with a Cu composition of 50 mol % or more of the entire subphase, and the crystal structure of the subphase is KHg 2 It is also stated that it is a type.

[0007] In Patent Document 5, R x Fe 100-x-y (V 1-a S a ) y (R is one or more rare earth elements including Y, x=5.5-18 atomic %, y=8-20 atomic %, a=0.05-0.7), and the main phase is ThMn 12 A rare earth permanent magnet with a body-centered tetragonal structure has been reported. This alloy consists of a main phase and a rare earth rich phase, and the RFe 2 It is described as being free of phases. [Prior art documents] [Patent documents]

[0008] [Patent Document 1] JP 2001-189206 A [Patent Document 2] International Publication No. 2017 / 164312 [Patent Document 3] JP 2017-112300 A [Patent Document 4] JP 2019-044259 A [Patent Document 5] Japanese Patent Application Publication No. 06-231920 Summary of the Invention [Problem to be solved by the invention]

[0009] As mentioned above, ThMn 12In order to obtain good magnetic properties in magnets having a main phase of this type compound, it has been proposed to form a structure consisting of a main phase and a grain boundary phase, as in Nd-Fe-B magnets, and non-magnetic phases such as La-rich phase (Patent Document 2) and R-Cu phase (Patent Documents 1 and 4) have been considered as the grain boundary phase. However, in reality, these phases segregate at grain boundary triple junctions and the like, making it difficult to form a grain boundary phase between two particles, and there has been a problem in that it is difficult to form a structure in which the surfaces of the main phase grains are covered by the grain boundary phase.

[0010] In addition, in Patent Document 3, Sm 5 Fe 17 Phase and SmCo 5 The surface of the main phase grains is surrounded by the ZnO phase, which pins the domain walls and improves the coercivity. 12 The grain surface of the Sm type compound 5 Fe 17 Phase and SmCo 5 It is difficult to achieve an organizational form that is encompassed by systemic aspects.

[0011] On the other hand, in Patent Document 5, ThMn 12 An alloy consisting of a main phase and an R-rich phase has been proposed. However, in reality, the composition range in which only two phases are formed in the R-Fe-V-Si quaternary system is extremely limited, so it is difficult to reproducibly produce this structure.

[0012] The present invention has been made in view of the above problems, and provides a method for producing a ThMn alloy having excellent magnetic properties. 12 The object of the present invention is to provide an anisotropic rare earth sintered magnet having a main phase of a compound of the type crystal. [Means for solving the problem]

[0013] As a result of extensive research to achieve the above object, the present inventors have discovered that ThMn 12 The inventors have discovered that an anisotropic rare earth sintered magnet having a main phase of a compound of this type crystal exhibits high coercivity when an interparticle grain boundary phase is formed between adjacent main phase grains, and have completed the present invention.

[0014] Accordingly, the present invention provides the following anisotropic rare earth sintered magnet and a method for producing the same. (1) The composition has the formula (R 1-a Zr a ) x (Fe 1-b Co b ) 100-x-y (M 1 1-c M 2 c ) y (R is one or more rare earth elements, Sm is essential, M 1 is one or more elements selected from the group consisting of V, Cr, Mn, Ni, Cu, Zn, Ga, Al, and Si; M 2 is one or more elements selected from the group consisting of Ti, Nb, Mo, Hf, Ta, and W, and x, y, a, b, and c are each represented by 7≦x≦15 atomic %, 4≦y≦20 atomic %, 0≦a≦0.2, 0≦b≦0.5, and 0≦c≦0.9, respectively; 12 The anisotropic rare earth sintered magnet comprises 80 volume % or more of a main phase made of a compound of a type crystal, the main phase has an average crystal grain size of 1 μm or more, and a grain boundary phase is formed between adjacent main phase grains. (2) The anisotropic rare earth sintered magnet according to (1), wherein the grain boundary phase between two particles contains 20 atomic % or more of R. (3) The anisotropic rare earth sintered magnet according to (1) or (2), wherein the grain boundary phase between 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 an R-rich phase is contained in the grain boundaries. (5)R(Fe,Co) 2 The anisotropic rare earth sintered magnet according to any one of (1) to (4), characterized in that the phase is contained in grain boundaries. (6) The R-rich phase and the R(Fe,Co) 2 % or more of the phases in total. (7) The Sm / R ratio inside the main phase grains is such that the Sm / R ratio inside the R-rich phase and the R(Fe,Co) 2The anisotropic rare earth sintered magnet according to any one of (4) to (6), wherein the Sm / R ratio of the first phase is lower than the Sm / R ratio of the second phase. (8) The anisotropic rare earth sintered magnet according to any one of (1) to (7), wherein the Sm / R ratio inside the main phase grain is lower than the Sm / R ratio in the outer periphery of the main phase grain. (9) The anisotropic rare earth sintered magnet according to (7) or (8), characterized in that the main phase grains do not contain Sm inside. (10) The anisotropic rare earth sintered magnet according to any one of (1) to (9), which exhibits a coercive force of 5 kOe or more at room temperature and has a temperature coefficient β of coercive force of -0.5% / K or more. (11)ThMn 12 The method for producing an anisotropic rare earth sintered magnet according to any one of (1) to (10), characterized in that an alloy containing a compound phase of a crystalline structure is pulverized, compacted in an applied magnetic field to form a green body, and then sintered at a temperature of 800°C or higher and 1400°C or lower. (12)ThMn 12 The method for producing an anisotropic rare earth sintered magnet according to (11) above, characterized in that an alloy containing a compound phase of the type crystal and an alloy having a higher R composition ratio and Sm / R ratio than that of the alloy are crushed and mixed, and then compacted under application of a magnetic field to form a green body. (13)ThMn 12 (11) or (12), characterized in that a material containing Sm is brought into contact with a sintered body having a compound phase of the type crystal as a main phase, and the Sm is diffused into the sintered body by heat-treating the sintered body at a temperature of 600°C or higher and lower than the sintering temperature. (14) The method for producing an anisotropic rare earth sintered magnet according to (13), wherein the material containing Sm that is brought into contact with the sintered body is at least one selected from the group consisting of Sm metal, an alloy containing Sm, a compound containing Sm, and vapor containing Sm, and is in the form of at least one selected from the group consisting of powder, thin film, ribbon, foil, and gas. (15) The method for producing an anisotropic rare earth sintered magnet according to any one of (11) to (14), further comprising subjecting the sintered body to a heat treatment at a temperature of 300 to 900°C. Effect of the Invention

[0015] According to the present invention, ThMn 12 In an anisotropic rare earth sintered magnet having a main phase of a compound of this type crystal, it is possible to obtain an anisotropic rare earth sintered magnet that exhibits excellent magnetic properties. [Brief description of the drawings]

[0016] [Figure 1] 1 is a HAADF image of the structure of the sintered magnet of Example 1 observed by STEM. [Diagram 2] 11 is a HAADF image of a different portion of the structure of the sintered magnet of Example 1 observed by STEM. [Diagram 3] 1 is a HAADF image of the structure of the sintered magnet of Comparative Example 1 observed with a STEM. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0017] The present invention relates to an anisotropic rare earth sintered magnet having 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 ThMn 12 The main phase is a compound of the type crystal, ThMn 12 The anisotropic sintered magnet contains 80% by volume or more of a main phase made of a compound of the type crystal, the main phase has an average crystal grain size of 1 μm or more, and two-particle grain boundary phases are formed between adjacent main phase grains. Note that 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. As the composition range is thus wide, the anisotropic rare earth sintered magnet of the present invention can be easily produced 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 is essential for Sm, and 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. R is the main phase of ThMn 12 It is an element necessary for forming a compound having a crystalline structure of the ThMn type. The content of R is set to 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 %, an α-Fe phase precipitates and sintering becomes difficult, while if it exceeds 15 atomic %, it becomes difficult to sinter the ThMn 12 The volume ratio of the ThMn type compound phase decreases, and good magnetic properties cannot be obtained. 12 When R is Sm, the anisotropy field H A Therefore, Sm is essential for the anisotropic rare earth sintered magnet of the present invention. When there is no difference in Sm concentration between the inside and the outer periphery of the main phase grains, the amount of Sm contained in R is preferably 5% or more of R in atomic ratio, more preferably 10% or more, and particularly preferably 20% or more. By having the Sm ratio in this range, H A The effect of increasing the magnetic field becomes sufficient to obtain a high coercive force.

[0019] On the other hand, since Sm is less produced than Y, La, Ce, Pr, Nd, etc. and is limited in terms of resources, it is preferable to use Sm as effectively as possible. Therefore, a structure in which Sm is concentrated in the outer periphery of the main phase grains may be used to obtain high coercivity with a smaller Sm content. In the case where the structure has different Sm concentrations in the inside and outer periphery of the main phase grains, the Sm contained in R is preferably 0.1 atomic % or more and 50 atomic % or less of R in atomic ratio. It is more preferable that the Sm content is 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 that R is a combination of one or more elements selected from Y, La, Ce, Pr, and Nd and Sm.

[0020] Zr, ThMn 12 It has the effect of substituting R in the type compound to improve phase stability. The amount of Zr substituting R should be 20% or less of R in terms of atomic ratio. If it exceeds 20%, it becomes a ThMn 12H type compounds A decreases, making it difficult to obtain high coercive force.

[0021] ThMn 12 It is known that in order for the crystalline structure to exist stably, a third element M is necessary in addition to the constituent elements R and Fe. 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 the third element. 1 is the third element, M 2 Compared to Fe, ThMn is an element that tends to form compounds with R more easily than Fe, or tends to bond less easily with both Fe and R. One of the features of the anisotropic rare earth sintered magnet of the present invention is that in the magnet structure, the main phase ThMn 12 Along with the type compounds, R-rich phases and R(Fe,Co) phases were observed at the grain boundaries. 2 The phase exists at the point where M exists as a third element. 1 By selecting the appropriate elements, it becomes easier to obtain a structure in which these three phases coexist stably. 1 and M. 2 If we combine these and write them as M, we get M 1 The atomic ratio of M is at least 10% or more, preferably 30% or more, and more preferably 50% or more. 1 If the content of M is less than 10%, the R-rich phase among the above three phases is not formed stably. 1 and M. 2 The total of M is 4 atomic % or more and 20 atomic % or less. When M is less than 4 atomic %, ThMn 12 If the content exceeds 20 atomic %, the amount of heterophase formed increases and good magnetic properties are not obtained. The R-rich phase is a phase with a higher concentration of rare earth elements than the main phase. 2 The phase is MgCu 2 This compound phase has a structure called a Laves phase.

[0022] M 2is one or more elements selected from Ti, Nb, Mo, Hf, Ta and W. 2 ThMn 12 It has the effect of stabilizing the M type crystal structure, but if it is contained in excess, 2 Carbides such as C phase and MgZn 2 (Fe,Co) type compound 2 M 2 Phases precipitate in the main phase and at grain boundaries. In particular, (Fe, Co) 2 M 2 The phase is, for example, Fe 2 Like the Ti phase, the composition may become Fe-richer than the stoichiometric composition, resulting in ferromagnetism, which can adversely affect the magnetic properties of the sintered magnet. 1 Not including M 2 If only M is selected, the R-rich phase is unlikely to be formed stably. 2 In the case of a composition containing M, its content is at least 90% or less of M in terms of atomic ratio.

[0023] The anisotropic rare earth sintered magnet of the present invention is 1 In addition, Fe is an essential constituent element. Furthermore, a part of Fe may be substituted with Co. The substitution with Co is 12 Curie temperature T of the compound c Increase the saturation magnetization M s The substitution rate of Co is set to 50% or less in atomic ratio. If the substitution rate exceeds 50%, M s On the other hand, the ratio of Fe and Co decreases. 1 and M 2 The balance is 3% by weight. In addition, it may contain unavoidable impurities that are taken in from raw materials or mixed in during the manufacturing process, specifically H, B, C, N, O, F, P, S, Mg, Cl, Ca, etc.

[0024] 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 ThMn 12 R(Fe,Co,M) type crystal structure 12It is preferable that elements such as C, N, and O, which are inevitably mixed in during the process of manufacturing sintered magnets, are not included in the main phase. However, if C, N, or O elements are detected in a composition analysis using an EPMA (electron probe microanalyzer) due to measurement variations, the method of preparing the observation sample, or the influence of the detection signals of other elements, it is considered that the H of the main phase A From the viewpoint of obtaining good results, the upper limit of each is preferably up to 1 atomic %. The average crystal grain size of the main phase is 1 μm or more, and preferably 1 μm to 30 μm. It is more preferable that it is in the range of 1.5 μm to 20 μm, and particularly preferably 2 μm to 10 μm. By setting the average crystal grain size in such a range, the residual magnetic flux density B due to the decrease in the degree of orientation of the crystal grains is reduced. r Decrease in coercive force H cJ The volume fraction of the main phase is good. r Or H cJ From the viewpoint of obtaining this, the content is 80% by volume or more, 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, based on the total volume of the magnet. The average crystal grain size of the main phase is a value measured as follows. The cross section of the sintered magnet was polished to a mirror finish, then immersed in an etching solution (such as a mixture of nitric acid + hydrochloric acid + glycerin) to selectively remove the grain boundary phase, and the cross section was observed at random at 10 or more points using a laser microscope. The cross-sectional area of ​​each particle was calculated from the obtained observation images using image analysis, and the average diameter when these were regarded as circles was taken as the average crystal grain size. The volume fraction of the main phase is a value measured as follows. Using EPMA, the structure of anisotropic rare earth sintered magnets was observed and the composition of each phase was analyzed. The main phase, R-rich phase, and R(Fe,Co) 2 The volume ratio of each phase was calculated as being equal to the area ratio in the backscattered electron image.

[0025] In order to effectively utilize Sm, a structure may be used in which Sm is concentrated in the outer periphery of the main phase grains, and grains with a lower Sm concentration exist inside the main phase grains. In this case, the thickness of the high Sm outer periphery is not particularly limited, but from the viewpoint of obtaining a sufficient effect of suppressing the nucleation of reverse magnetic domains in the outer periphery of the main phase grains and preventing the Sm content of the entire sintered body from becoming too high, the thickness is preferably 1 nm to 2 μm, and particularly preferably 2 nm to 1 μm. This type of morphology is also known as an R-rich phase or an R(Fe,Co) 2 This is achieved by increasing the Sm / R ratio (atomic ratio of Sm to R) in the phase compared to the Sm / R ratio inside the main phase grains. It is more preferable if the main phase grains do not contain Sm inside. It is also acceptable for some main phase grains to contain a uniform Sm concentration distribution.

[0026] R-rich phase and R(Fe,Co) 2 The R-rich phase is formed at the grain boundary of the magnet structure. The grain boundary includes grain boundary triple junctions as well as grain boundary phases between two grains. Here, the R-rich phase is defined as a phase containing 40 atomic % or more of R. 1 When the above composition is used, including the elements, the main phase, R(Fe,Co) 2 It has been found that it is easy to obtain a magnet containing three phases: M phase, R-rich phase, and R-rich phase. 1 In the case of Sm-Fe-Ti ternary sintered magnets that do not contain any elements, Sm(Fe,Ti) 12 Main phase and SmFe 2 , Fe 2 There is a composition region where the three phases of Ti (excluding oxides, etc.) are in equilibrium, but Sm(Fe,Ti) 12 Since the main phase and the Sm-rich phase are difficult to equilibrate at low temperatures below 400°C, the Sm-rich phase does not form as a stable phase. 1 In the case of the Sm-Fe-V ternary system, where one of the elements is V, no Fe-V binary compound is formed, but instead a Sm-rich phase with a high Sm concentration is formed, resulting in Sm(Fe,V) 12 , SmFe 2 A magnet containing three phases, namely, a Sm-rich phase and a Sm-rich phase, can be obtained. 1 , M 2In the Sm-Fe-V-Ti quaternary system, which contains both Sm(Fe,V,Ti) 12 , Fe 2 (V,Ti), SmFe 2 Based on this knowledge, the anisotropic rare earth sintered magnet of the present invention has an R-rich phase and an R(Fe,Co) phase at the grain boundaries. 2 To form the phase, a given amount of M 1 A composition is selected that includes the elements.

[0027] R-rich phase and R(Fe,Co) 2 The R-rich phase has four main effects. The first effect is to promote sintering. At the sintering temperature, both the R-rich phase and the R(Fe,Co) 2 The R-rich phase and the R(Fe,Co) phase are also melted and become liquid phase, so liquid phase sintering proceeds and sintering is completed more quickly than solid phase sintering without these phases. 2 When these phases coexist, the liquid phase formation temperature tends to be lower than when only one of the phases is present, and liquid phase sintering proceeds more rapidly.

[0028] The second effect is the cleaning of the main phase grain surfaces. Since the anisotropic rare earth sintered magnet of the present invention has a nucleation-type coercive force mechanism, it is desirable for the main phase grain surfaces to be smooth so that nucleation of reverse magnetic domains is less likely to occur. 2 The phase is transformed into ThMn during the sintering process or the subsequent aging process. 12 This cleaning effect smooths the surface of the crystal grains of R(Fe,Co) type compounds, suppressing the nucleation of reverse magnetic domains, which is a factor in reducing coercivity. 2 The phase may be any other phase with R less than 40 atomic %, e.g., RM 3 , R.M. 2 , R(Fe,Co)M and R(Fe,Co) 2 M 2 Compared with compound phases such as ThMn 12 It has a relatively high wettability with respect to the main phase and is easy to cover the surfaces of the main phase particles, so it has a large cleaning effect.

[0029] The third effect is the formation of grain boundary phases between two grains. In magnets containing R-rich phases in the structure, the adjacent ThMn phases can be formed by optimal sintering or aging treatment. 12 Between the main phase grains of the ThMn type compound, a grain boundary phase containing more R than the main phase is formed. This weakens the magnetic interaction between the main phase grains, and the sintered magnet exhibits a high coercive force. 12 The compositional region in which only the two phases, the main phase of the ThMn-type compound and the R-rich phase, are in equilibrium is extremely limited, and considering the compositional variation, it is difficult to stably manufacture such magnets. 12 Main phase of type compound, R-rich phase and R(Fe,Co) 2 By making the magnet contain three phases, it is possible to stably form a structure in which the surfaces of the main phase grains are covered by the grain boundary phase between two particles. In addition, in magnets that do not contain an R-rich phase, it is difficult to form the grain boundary phase between two particles, or it is difficult for the grain boundary phase between two particles to cover the surfaces of the main phase grains, making it difficult to obtain a magnet that exhibits sufficient coercivity.

[0030] The fourth effect is to increase the Sm concentration in the grain boundaries. When the grain boundary diffusion method is applied as a manufacturing method to make the Sm concentration different between the inside and the outer periphery of the main phase grains, the R-rich phase and the R(Fe,Co) phase existing in the grain boundaries are mixed. 2 The R-rich phase or R(Fe,Co) becomes liquid during the diffusion process and plays a role in diffusing and penetrating the Sm placed on the sintered body into the interior. 2 The Sm / R ratio in at least one of the phases is higher than the Sm / R ratio in the main phase grains. 12 By using an alloy mainly composed of a type compound phase and an alloy with a higher R composition ratio and Sm / R ratio than that of the type compound phase, the R-rich phase or R(Fe,Co) 2 The Sm / R ratio in at least one of the phases is higher than the Sm / R ratio in the main phase grains. 2 As Sm becomes concentrated in the grain boundary phase, the Sm concentration in the outer crust of the main phase grains that contact these grain boundary phases also increases, and the H A This improves the coercive force of the sintered magnet.

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

[0032] On the other hand, R(Fe,Co) 2 The phase is MgCu 2 The Laves compound is a type crystal, but when the composition is analyzed using an EPMA or the like, taking into consideration measurement variations, it is assumed that the R content is 20 atomic % or more and less than 40 atomic %. In addition, part of Fe and Co may be replaced by the M element. However, the amount of M to be replaced is MgCu 2 The range is set so that the type crystal structure is maintained.

[0033] R(Fe,Co) in the anisotropic rare earth sintered magnet of the present invention 2 The magnetic phase is a ferromagnetic or ferrimagnetic phase with a Curie temperature T c RFe is the phase where the temperature is above room temperature (23°C). 2 is CeFe 2 Except for T c is above room temperature, and CeFe 2 If 10% or more of R is replaced by other elements, T c On the other hand, RCo 2 GdCo 2 Except for T c is below room temperature or in the paramagnetic phase, but in the anisotropic rare earth sintered magnet of the present invention, the atomic ratio of Fe substituted by Co is 0.5 or less, so in most cases R(Fe,Co) 2Generally, soft magnetic phases contained in the structure often have a negative effect on magnetic properties, but in the anisotropic rare earth sintered magnet of the present invention, the R(Fe,Co) 2 It is considered that the cleaning effect of the phase on the main phase grain surface and the effect of forming a grain boundary phase between two particles are greater, and that even the magnetic phase contributes to an increase in coercivity.

[0034] R-rich phase and R(Fe,Co) 2 The total amount of the formed phases is preferably 1% by volume or more, more preferably 1% by volume or more and less than 20% by volume. Also, the total amount of the formed phases is preferably 1.5% by volume or more and less than 15% by volume, and even more preferably 2% by volume or more and less than 10% by volume. By setting the total amount of the formed phases in such a range, the area of ​​contact with the main phase grains is secured, and the H cJ It is easy to achieve the effect of increasing the size. r The decrease in magnetic properties is also suppressed, making it easier to obtain the desired magnetic properties.

[0035] As described above, in the anisotropic rare earth sintered magnet of the present invention, the R-rich phase and the R(Fe,Co) phase are present at the grain boundaries. 2 The phase exists, and ThMn 12 The grain boundary phase between two grains is formed between adjacent main phase grains made of the type compound. The surface of the main phase grain is covered by the grain boundary phase between two grains, which weakens the magnetic interaction between the main phase grains and results in a high coercive force.

[0036] The grain boundary phase between the two grains may be amorphous with a disordered atomic arrangement, or may have a regular atomic arrangement. 2It may be the same phase as the phase. When the grain boundary phase between two particles is observed using a device such as a STEM (scanning transmission electron microscope), the composition preferably contains 20 atomic % or more of R. By setting the thickness within this range, the magnetic coupling between the main phase grains can be sufficiently reduced, and high coercivity can be easily obtained. In addition, the thickness of the grain boundary phase between two particles is preferably 0.5 nm or more. This makes it easy to ensure the magnetic isolation effect between the main phase grains, and a sufficient coercivity improvement effect can be obtained. Furthermore, the thickness is preferably 1 μm or less, more preferably 0.5 μm or less, and even more preferably 100 nm or less. If the thickness is within this range, it is easy to suppress the effect of the decrease in magnetic properties due to the decrease in the volume ratio of the main phase grains from being greater than the effect of the increase in coercivity. The thickness of the grain boundary phase between two grains was measured from the STEM image as follows. Using a STEM device (JEM-ARM200F manufactured by JEOL Ltd.), at least three or more locations on one sample were observed where adjacent main phase grains were in contact with each other. The thickness of the grain boundary phase between two grains was measured from the HAADF (High-Angle Annular Dark Field) images obtained by the observation, and the average value of these thicknesses was taken as the thickness of the grain boundary phase between two grains.

[0037] In addition, the anisotropic rare earth sintered magnet of the present invention may contain R oxides, R carbides, R nitrides, 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, more preferably 5 volume % or less, and particularly preferably 3 volume % or less.

[0038] It is preferable to have as few phases other than those mentioned above as possible. For example, R 2 (Fe,Co,M) 17 Phase, R 3 (Fe,Co,M) 29 If the phases are present in the magnet structure, their formation amount should be less than 1 volume % each in order to suppress the effect on magnetic properties and the resulting decrease in coercivity. Also, in order to ensure a sufficient proportion of the main phase, (Fe, Co) 2 RM with M phase or R less than 40 atomic % 3 , R.M.2 , R(Fe,Co)M, R(Fe,Co) 2 M 2 etc. are each preferably less than 1 volume %. The total amount of these phases is preferably 3 volume % or less. Furthermore, from the viewpoint of preventing a significant decrease in magnetic properties, it is preferable that the anisotropic rare earth sintered magnet of the present invention does not contain the α-(Fe,Co) phase.

[0039] Next, the manufacturing method will be described. The anisotropic rare earth sintered magnet of the present invention is manufactured by powder metallurgy. First, to prepare the raw alloy, metal raw materials of R, Fe, Co, M, alloys, ferroalloys, etc. are used, and the composition of the sintered body finally obtained is adjusted to a predetermined value, taking into consideration the loss of raw materials during the manufacturing process. These raw materials are melted in a high-frequency furnace or an arc furnace to prepare the alloy. The molten metal may be cooled 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 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 magnetic field molding process, resulting in B. r If α-Fe precipitates in the alloy, it is necessary to remove the α-Fe and 12 The alloy may be subjected to a heat treatment so as to increase the amount of the type compound phase formed. The alloy may be an alloy of a single composition, or may be prepared by preparing a plurality of alloys of different compositions and mixing the powders in a post-process.

[0040] The above-mentioned raw alloy is coarsely pulverized into powder having an average particle size of 0.05 to 3 mm by mechanical pulverization such as a Braun mill or hydrogen pulverization. Alternatively, the HDDR method (hydrogen disproportionation desorption recombination method) used as a manufacturing method for Nd-Fe-B magnets may be applied. The coarse powder is then finely pulverized by a ball mill or a jet mill using high-pressure nitrogen to obtain powder having an average particle size of 0.5 to 20 μm, more preferably 1 to 10 μm. Before or after the fine pulverization process, a lubricant or the like may be added as necessary. Next, the alloy powder is molded using a magnetic field press while the easy axis of magnetization of the alloy powder is aligned in an applied magnetic field to obtain a powder compact. The molding is preferably performed in a vacuum, a nitrogen gas atmosphere, an inert gas atmosphere such as Ar, or the like to suppress oxidation of the alloy powder.

[0041] The step of sintering the powder compact is carried out in a sintering furnace in a vacuum or inert atmosphere at a temperature of 800°C to 1400°C. At temperatures below 800°C, sintering does not proceed sufficiently and a high sintered density cannot be obtained. At temperatures above 1400°C, the sintering temperature is too high, and the sintered density is too low. 12 The main phase of the α-Fe-type compound decomposes and α-Fe precipitates. The sintering temperature is 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. 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 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 it can be cooled to 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 coercive force, an aging heat treatment at 300 to 900°C for 0.5 to 50 hours may be further performed. By optimizing the sintering and aging conditions according to the composition, powder grain size, etc., it is possible to obtain a high-temperature sintered product with a high coercive force. cJ The sintered body is then cut and ground into a specified shape, and magnetized to produce a sintered magnet.

[0042] On the other hand, the Sm / R ratio inside the main phase grains is R-rich phase and R(Fe,Co) 2 Examples of methods for producing anisotropic rare earth sintered magnets in which main phase grains exist that have a lower Sm / R ratio than the phase include the two-alloy method and the grain boundary diffusion method.

[0043] When using the two-alloy method, two types of raw material alloys with different compositions are prepared using metal raw materials, alloys, ferroalloys, etc. of R, Fe, Co, and M. Three or more types of alloys may be used. In this case, ThMn 12 It is preferable to combine alloy A, which is mainly composed of a type compound phase and has a relatively low Sm / R ratio, with alloy B, which has a relatively higher R composition ratio and Sm / R ratio, and adjust the average composition to a predetermined composition. These alloys are produced by casting or strip casting, and then pulverized. The process of mixing the alloy powders may be performed in a coarse powder state before fine pulverization, or after fine pulverization. The mixture is then molded and sintered to produce a sintered body. Aging heat treatment may be performed to improve the coercive force.

[0044] In the two-alloy sintered magnet, the composition of alloy A is mainly responsible for the formation of ThMn 12 The main phase is composed of a type compound, and the R-rich phase, R(Fe,Co) is mainly formed due to the composition of alloy B. 2 As a result, the R-rich phase and R(Fe,Co) phase formed at the grain boundaries are 2 The Sm / R atomic ratio of the grain boundary phase is higher than that of the interior of the main phase grains. In addition, some of the Sm in the grain boundary phase substitutes for R atoms in the surface layer of the main phase grains, forming a core-shell structure with different Sm concentrations between the grain surface and the interior, thereby increasing the coercive force.

[0045] 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 in the same manner as described above. At this time, R in the sintered body composition may or may not contain Sm.

[0046] Next, the obtained sintered body is subjected to grain boundary diffusion of Sm. After cutting and grinding the sintered body as necessary, a diffusion material selected from compounds containing Sm, such as metals, alloys, oxides, fluorides, oxyfluorides, hydrides, and carbides, is placed on the surface in the form of powder, thin film, ribbon, foil, or the like. For example, a 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 vapor deposition, sputtering, CVD, or the like. The amount of the material to be placed is 10 to 1000 μg / mm 2 It is preferable that the concentration is 20 to 500 μg / mm 2 In this range, H cJ In addition, it is possible to suppress the increase in manufacturing cost due to the high Sm content. In addition, by utilizing the property of high vapor pressure of Sm, Sm metal or Sm alloy may be heat-treated together with the sintered compact in the same chamber, and the Sm vapor may be brought into contact with the sintered compact.

[0047] The sintered body, with Sm placed on the surface, is heat-treated in a vacuum or in an inert gas atmosphere. The heat treatment temperature is preferably 600°C or higher and lower than the sintering temperature, and more preferably 700°C or higher and lower than 1100°C. The heat treatment time is preferably 0.5 to 50 hours, and more 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 more preferably 2 to 10°C / min. In order to improve the coercive force, an aging heat treatment may be further performed at 300 to 900°C for 0.5 to 50 hours.

[0048] The Sm placed on the sintered body turns into R-rich phase or R(Fe,Co) by heat treatment. 2 As the Sm concentration in the grain boundary phase increases, Sm is substituted for R atoms in the surface layer of the main phase grains that contact the grain boundary phase, and the Sm / R ratio in the surface layer of the main phase grains becomes higher than the Sm / R ratio in the interior of the main phase grains, resulting in the formation of H cJ increases.

[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 temperature.r and a coercive force H of at least 5 kOe cJ Shows room temperature H cJ It is more preferable that the temperature coefficient of coercivity β is -0.5% / K or more. Here, β=ΔH cJ / ΔT×100 / H cJ (20℃)(ΔH cJ =H cJ (20℃)-H cJ (140°C, ΔT=20-140(°C)). The anisotropic rare earth sintered magnet of the present invention has a smaller change in coercive force with temperature compared to Nd-Fe-B sintered magnets, and is therefore suitable for use at high temperatures. EXAMPLES

[0050] EXAMPLES The present invention will be specifically described below with reference to examples and comparative examples, 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 the alloy was melted in an Ar gas atmosphere in a high-frequency induction furnace to produce a cast alloy. In order to eliminate the primary crystal α-Fe, the alloy was heat-treated at 900°C for 50 hours. The structure of the obtained alloy was observed using a laser microscope (Olympus Corporation, LEXT OLS4000), and the observed images confirmed that the average crystal grain size of the main phase was 5 μm or more. The alloy was subjected to a hydrogen absorption treatment and a dehydrogenation treatment by heating in a vacuum at 400°C to produce a coarse powder, which was then pulverized in a jet mill in a nitrogen gas flow to produce a fine powder with an average grain size of 1.8 μm. The fine powder was then filled into a mold of a molding device in an inert gas atmosphere, and oriented in a magnetic field of 15 kOe (=1.19 MA / m) while being subjected to a 0.6 Ton / cm in the direction perpendicular to the magnetic field. 2 This powder compact was sintered in an Ar gas atmosphere at 1140°C for 3 hours, and then 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 spectrometry analyzer (Hitachi High-Tech Science Corporation, SPS3520UV-DD) was Sm 10.2 Fe bal. V 14.9 Al 0.5 S 0.2 In addition, X-ray diffraction revealed that ThMn 12 The sintered body was observed for its structure and the composition of the formed phase was analyzed using an EPMA device (JXA-8500F, manufactured by JEOL Ltd.), and the R-rich phase and R(Fe,Co) phase were found at the grain boundary triple junction. 2 It was confirmed that the main phase, the R-rich phase and the R(Fe,Co) phase were formed. 2 The volume ratio of the phase was calculated as being equal to the area ratio in the backscattered electron image. As a result, the grain boundary of the sintered body structure was found to contain an R-rich phase and an R(Fe,Co) 2 The phases each existed at 1% by volume or more. The average crystal grain size of the main phase calculated from the results of etching and observing the sintered sample was 9.9 μm. Furthermore, the room temperature H cJ is 8.5kOe, and H cJ The temperature coefficient β was -0.46% / K.

[0053] From this sintered body, a thin specimen for observation was cut out using a FIB-SEM device (Scios Dual Beam, manufactured by FEI) and observed using a STEM device (JEM-ARM200F, manufactured by JEOL Ltd.). The obtained HAADF (High-Angle Annular Dark Field) images are shown in Figures 1 and 2. From Figure 1, it can be confirmed that an interparticle grain boundary phase 13 exists in the grain boundary portion sandwiched between two main phase grains 11 and 12. The thickness of the interparticle grain boundary phase 13 in this case was approximately 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 S 1.1 In addition, in Fig. 2, which shows another portion of the same sintered body, the grain boundary phase 13 between two grains is about 1.4 nm thick and has a composition of Sm31.9 Fe bal. V 9.5 Al 2.7 S 0.6 The results are shown in Tables 1 to 3 and 5.

[0054] [Comparative Example 1] The composition was adjusted using Sm metal, electrolytic iron, and Ti metal, and a cast alloy was produced in a high-frequency induction furnace in the same manner as in Example 1, and further heat-treated at 900°C for 50 hours. The structure of the obtained alloy was observed using a laser microscope, and it was confirmed from the observed image that the average crystal grain size of the main phase was 5 μm or more. Crushing and compacting in a magnetic field were performed in the same manner as in Example 1, and the alloy was sintered in an Ar gas atmosphere at 1175°C for 3 hours, and then cooled to room temperature at a cooling rate of 13°C / min to obtain a sintered body of Comparative Example 1. The composition value of the sintered body analyzed by ICP method was Sm 9.7 Fe bal. Ti 8.1 In addition, the X-ray diffraction measurement revealed that the main phase of Comparative Example 1 was ThMn 12 The formation phase was examined with EPMA and found to be R(Fe,Co) 2 Although the phase was present, no R-rich phase was formed, and a fine TiC phase was precipitated. When the magnetic properties were measured using a BH tracer, this Comparative Example 1 showed a low coercive force of only 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, no intergranular grain phase was formed between the two grains as seen in Example 1. The results are shown in Tables 1 to 3.

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

[0056] [Comparative Examples 2 to 6] Sintered body samples of Comparative Examples 2 to 5 were produced 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 the sintering was not sufficient, and 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 %, and an R-rich phase was not observed, and an RFeSi phase of PbClF type crystal was formed. In Comparative Example 5, KHg 2 Rcu type crystal 2 In Comparative Example 6, the total content of M is less than 4 atomic %, and the structure contains no ThMn phase. 12 No Th type crystals were observed. 2 Zinc 17 The main phase of the type crystal was formed.

[0057] [Comparative Example 7] The composition was adjusted using Sm metal, electrolytic iron, Ti metal, and V metal, and the raw material molten metal was cooled on a Cu roll rotating at a peripheral speed of 20 m / sec to produce a quenched ribbon raw material alloy. The ribbon had a thickness of 10 to 50 μm, and the alloy structure obtained was observed using a laser microscope. From the observed images, it was confirmed that the average crystal grain size was at least smaller than 1 μm, although it was too fine to measure. This alloy ribbon was pulverized in a ball mill, and then powder with a grain size of 300 μm or less was selected using a sieve and hot pressed at 750°C in an Ar atmosphere. The average crystal grain size of the main phase grains was as small as 0.2 to 0.3 μm, and the composition of the main phase and grain boundary phase could not be identified by EPMA. In addition, the easy axis of magnetization of the main phase was not aligned, so the low B r 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, and melted in an Ar gas atmosphere using a high-frequency induction furnace, and then strip-cast on a water-cooled Cu roll to produce a quenched ribbon alloy with a thickness of about 0.2 to 0.4 mm and a composition of Nd 7.5 atomic %, Y 1.0 atomic %, Si 13.0 atomic %, Hf 1.0 atomic %, and the balance Fe. The average crystal grain size in the minor axis direction of the alloy obtained from an image observed with a laser microscope was 2.5 μm. This alloy was subjected to hydrogen absorption treatment at room temperature, and then subjected to dehydrogenation treatment by heating at 400 ° C in a vacuum to obtain a coarse powder (this is referred to as the actual 9A powder). On the other hand, an alloy ingot with a composition of Sm 40 atomic %, Ga 10 atomic %, Cu 5 atomic %, and the balance Co was produced using Sm metal and electrolytic iron as raw materials using a high-frequency induction furnace, and was then mechanically crushed to obtain a coarse powder (referred to as the actual 9B powder). The actual 9A powder and the actual 9B powder were mixed in a weight ratio of 95:5, and then pulverized in a jet mill in a nitrogen stream to produce a fine powder having an average particle size of 1.8 μm.

[0059] This mixed powder was subjected to magnetic field molding in the same manner as in Example 1, sintered in an Ar gas atmosphere at 1200°C for 3 hours, cooled to room temperature at a cooling rate of 12°C / min, and further heat-treated in an Ar gas atmosphere at 650°C for 1 hour to obtain a sintered body of Example 9. The composition of the sintered body sample was Sm 1.8 Nd 7.2 Y 1.0 Fe bal. Co 1.0 S 12.8 Ga 0.6 Cu 0.4 Hf 1.0 In addition, X-ray diffraction measurements showed that the main phase of this sintered body was ThMn 12 It was confirmed that the crystals were of the above type. The composition of the main phase measured by EPMA was Nd 6.4 Y 1.1 Fe bal. Co 1.0 S 12.7 Ga 0.5 Cu 0.1 Hf 1.1 does not contain Sm, but the outer shell of the grains does not contain Sm 3.5 Nd 3.0 Y 1.0 Fe bal. Co1.0 S 13.0 Ga 0.4 Cu 0.1 Hf 0.9 It was confirmed that the Sm / R ratio inside the grains was lower than that of the surface layer. In addition, the structure of the sintered body was observed by EPMA and the composition of each phase was analyzed, and it was found that the R-rich phase and the R(Fe,Co) phase were present at the grain boundaries. 2 It was confirmed that the phase was present at 1 volume % or more. 2 The phase was confirmed. 2 (Fe,Co,M) 17 Phase, R 3 (Fe,Co,M) 29 No α-Fe or α-Fe phase was observed. Note that the total of the phase ratios is less than 100% because oxides and other phases are also present.

[0060] R-rich phase, R(Fe,Co) 2 Phase and RCu 2 The compositional analysis values ​​of the phases are Sm 26.8 Nd 31.7 Y 0.1 Fe bal. S 36.2 Ga 5.2 , Sm 17.2 Nd 17.4 Y 0.2 Fe bal. Co 0.4 S 0.3 Hf 0.1 , Sm 15.9 Nd 18.6 Fe bal. Cu 65.2 From this, the Sm / R ratio in the grain interior is the R-rich phase and the R(Fe,Co) 2 It was confirmed that the Sm / R ratio of the main phase was lower than that of the Fe,Co 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. R(Fe,Co) 2 The Curie temperature T of alloys of the same composition prepared based on the phase analysis values c was 318°C. For Example 9, STEM observation was performed in the same manner as in Example 1, and it was confirmed that an intergranular phase was present in the grain boundary portion between two main phase grains. The measured composition of the intergranular phase was Sm21.7 Nd 24.5 Fe bal. Co 0.5 S 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, and the composition was melted in an Ar gas atmosphere using a high-frequency induction furnace, and then strip-cast on a water-cooled Cu roll to produce an alloy strip with a thickness of about 0.2 to 0.4 mm. The average grain boundary phase interval of this alloy was calculated to be 4.1 μm. The alloy was subjected to the same hydrogen absorption and dehydrogenation treatment as in Example 9 to obtain coarse powder, which was then pulverized by 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, cooled to room temperature at a cooling rate of 11°C / min, and removed 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, and once removed from the furnace, was further subjected to aging treatment at 520°C for 2 hours to obtain Example 10.

[0062] The sintered body sample of Example 10 was subjected to ICP analysis, and the composition was Sm 2.4 Ce 7.7 La 1.1 Fe bal. Co 0.6 S 12.6 Mo 0.9 From the X-ray diffraction measurement of the powder obtained by crushing a part of the sample, the crystal structure of the main phase was found to be ThMn 12 In addition, the structure of the sintered body was observed by EPMA and the composition of each phase was analyzed, and it was found that the R-rich phase and the R(Fe,Co) phase were present at the grain boundaries. 2 It was confirmed that the R phase was present at 1% by volume or more. 2 (Fe,Co,M) 17 Phase, R 3 (Fe,Co,M) 29 No α-Fe or α-Fe phase was observed. Note that the total of the phase ratios is less than 100% because oxides and other phases are also present.

[0063] The compositional analysis values ​​of the center and outer periphery of the main phase grains by EPMA were Ce, 7.3 La 0.1 Fe bal. Co 0.5 S 13.5 Mo 1.0 , Sm 3.2 Ce 4.2 La 0.3 Fe bal. Co 0.5 S 13.5 Mo 1.0 It was confirmed that the Sm / R ratio in the grain interior was lower than that in the surface layer. 2 The compositional analysis values ​​of the phases are Sm 31.1 Ce 19.6 La 23.2 Fe bal. Co 2.7 S 0.6 Mo 0.2 , Sm 15.0 Ce 18.9 La 0.5 Fe bal. Co 0.2 S 0.7 Mo 0.2 No Sm was detected inside the main phase grains, whereas the R-rich phase and R(Fe,Co) in the grain boundaries were detected. 2 It was confirmed that the phase contains Sm and has a high Sm / R ratio. R(Fe,Co) 2 Based on the phase analysis, an alloy with the same composition was prepared by arc melting and homogenized at 800°C for 20 hours. After that, magnetization-temperature measurements were performed using a VSM. The Curie temperature T c The average crystal grain size of the main phase calculated from the results of etching and observing the sintered body of Example 9 was 12.3 μm. Furthermore, when the magnetic properties were measured with a BH tracer, the room temperature coercivity H cJ The value was 6.3 kOe. cJ The temperature coefficient β was -0.48% / K.

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

[0065] [Comparative Example 8] The sintered body of Comparative Example 9 was produced in the same manner as the sintered body of Example 10, except that the step of heat treating the sintered body together with Sm metal was not carried out and aging treatment was carried out at 520° C. for 2 hours.

[0066] The composition of the sintered body of Comparative Example 8 is Ce without Sm. 8.3 La 1.3 Fe bal. Co 0.6 S 13.0 Mo 0.9 The composition analysis value of the main phase grains is Ce. 7.5 La 0.3 Fe bal. Co 0.6 S 13.1 Mo 0.9 At the grain boundaries, R(Fe,Co) 2 No phase exists and the composition is Ce 33.1 La 29.6 Fe bal. S 37.3 and Ce 23.3 La 54.3 Fe bal. Co 0.8 S 0.6 Mo 0.1 Two types of R-rich phases were observed. cJ 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 grains

Claims

1. The composition has the formula (R １－ａ Zr ａ ) ｘ (Fe １－ｂ Co ｂ ) １００－ｘ－ｙ (M １ １－ｃ M ２ ｃ ) ｙ (R is 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; M １ represents one or more elements selected from the group consisting of V, Cr, Mn, Ni, Cu, Zn, Ga, Al, and Si; M ２ is one or more elements selected from the group consisting of Ti, Nb, Mo, Hf, Ta, and W, and x, y, a, b, and c are each represented by 7≦x≦15 atomic %, 4≦y≦20 atomic %, 0≦a≦0.2, 0≦b≦0.5, and 0≦c≦0.9, respectively; １２ The main phase is composed of a compound of the Fe-type crystal in an amount of 80% by volume or more, and the R-rich phase and the R(Fe, Co) ２ The main phase has an average crystal grain size of 1 μm or more, and an intergranular grain phase is formed between adjacent main phase grains. The Sm / R ratio in the inside of the main phase grain is ２ The anisotropic rare earth sintered magnet is characterized in that the Sm / R ratio of the magnet is lower than that of the magnet.

2. 2. The anisotropic rare earth sintered magnet according to claim 1, wherein the grain boundary phase between two particles contains 20 atomic % or more of R.

3. 3. The anisotropic rare earth sintered magnet according to claim 1, wherein the grain boundary phase between two particles has a thickness of 0.5 nm or more.

4. The R-rich phase and the R(Fe, Co) ２ 4. The anisotropic rare earth sintered magnet according to claim 1, wherein the phases are contained in an amount of 1 volume % or more in total.

5. 5. The anisotropic rare earth sintered magnet according to claim 1, wherein the Sm / R ratio inside the main phase grain is lower than the Sm / R ratio in the outer periphery of the main phase grain.

6. The anisotropic rare earth sintered magnet according to any one of claims 1 to 5, characterized in that the main phase grains do not contain Sm inside.

7. 7. The anisotropic rare earth sintered magnet according to claim 1, which exhibits a coercive force of 5 kOe or more at room temperature and a temperature coefficient β of coercive force of -0.5% / K or more.

8. ThMn １２ The method for producing an anisotropic rare earth sintered magnet according to any one of claims 1 to 7, characterized in that an alloy containing a compound phase of the type crystal is pulverized, compacted under application of a magnetic field to form a green body, and then sintered at a temperature of 800°C or higher and 1400°C or lower.

9. ThMn １２ 9. The method for producing an anisotropic rare earth sintered magnet according to claim 8, characterized in that an alloy containing a compound phase of the type crystal and an alloy having a higher R composition ratio and Sm / R ratio than those of the type crystal are pulverized and mixed, and then compacted under application of a magnetic field to form a green body.

10. ThMn １２ 10. The method for producing an anisotropic rare earth sintered magnet according to claim 8, wherein a material containing Sm is brought into contact with a sintered body having a compound phase of the type crystal as a main phase, and the material is heat-treated at a temperature of 600°C or higher and lower than the sintering temperature, thereby diffusing Sm into the sintered body.

11. 11. The method for producing an anisotropic rare earth sintered magnet according to claim 10, characterized in that the material containing Sm that is brought into contact with the sintered body is one or more selected from the group consisting of Sm metal, an Sm-containing alloy, a compound containing Sm, and vapor containing Sm, and is in the form of one or more selected from the group consisting of powder, thin film, ribbon, foil, and gas.

12. The method for producing an anisotropic rare earth sintered magnet according to any one of claims 8 to 11, characterized in that the sintered body is subjected to a heat treatment at a temperature of 300 to 900°C.

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