WO2016162990A1 - Rare earth permanent magnet and method for producing same - Google Patents

Abstract

When a non-equilibrium ThMn₁₂-type intermetallic compound formed by rapid quenching is converted into a magnet, no low melting liquid phase is formed so that there is a problem in densification. In the present invention, a liquid phase with a melting point of about 820^°C can be formed by adding a trace amount of Cu to a quaternary system comprising Sm-Y-Fe-Co as main components and thus densification can be promoted in the process of producing a bulk magnet.

Description

Rare earth permanent magnet and manufacturing method thereof

The present invention relates to a ferromagnetic alloy and a method for producing the same.

In recent years, development of magnets with a reduced content of rare earth elements has been demanded. The rare earth element in this specification is at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanoid. Here, the lanthanoid is a general term for 15 elements from lanthanum to lutetium.

As a ferromagnetic alloy having a relatively small composition ratio of rare earth elements contained, RFe ₁₂ (R is at least one kind of rare earth elements) having a body-centered tetragonal ThMn ₁₂ type crystal structure is known. However, RFe ₁₂ has a problem that the crystal structure is thermally unstable in the binary system. Patent Document 1 teaches that a ThMn ₁₂ type is produced in a Y—Fe binary system by selecting Y as R and using a superquenching method. In Patent Document 2, by adding at least one element T selected from Cu, Ag, Bi, Mg, Sn, Pb and In, a phase having a lower melting point and a non-magnetic property compared to the main phase of ThMn ₁₂ type To form a two-phase structure with at least the main phase.

JP 2014-47366 A JP 2001-189206 A

In the ferromagnetic alloy of Patent Document 1, a part of the Fe element is replaced with a structural stabilizing element M (M = Si, Al, Ti, V, Cr, Mn, Mo, W, Re, Be, Nb, etc.). It has high magnetization because it does not exist, but the magnitude of magnetization and magnetic anisotropy is still small for practical use, and in the binary system, as shown in the phase diagram, it becomes dense, high coercive force, high mechanical strength The low-melting-point grain boundary phase necessary for this is not generated. Further, in Patent Document 2, by adding a T element to the ThMn ₁₂ type, a liquid phase having a melting point lower than that of the main phase is generated, and a high coercive force is generated by functioning as a grain boundary phase. Only applies to

Thus, in the conventional method, there is a problem in terms of densification in order to magnetize the non-equilibrium ThMn ₁₂ type intermetallic compound produced by the ultra-quenching method.

One aspect of the present invention for solving the above problems is that R′-LRE (Light Rare-earth) —Fe—Co based ferromagnetic alloy (R ′ is at least one selected from Y and Gd, and LRE is La , Ce, Nd, Pr, and Sm having a composition of the formula R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (where 0 <x <0.5, 0 <y <0.5, 10 <z <19, 0.01 ≦ α <0.5), and the main phase has an intermediate crystal structure between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure. It is a rare earth permanent magnet characterized by being an R′-LRE—Fe—Co based ferromagnetic compound.

In a more preferred specific example, the range of 0.1 ≦ α ≦ 0.4 is specified. Further, the density is preferably higher than that when α is 0.

Another aspect of the present invention for solving the above problems is an R′—LRE—Fe—Co based ferromagnetic alloy (R ′ is at least one selected from Y and Gd, LRE is La, Ce, Nd, The composition of at least one selected from Pr and Sm is represented by the formula R ′ _1-x LRE _{x + β} (Fe _1-y Co _y ) _z Cu _vβ (0 <x <0.5 and 0 <y <0.5, 10 <z <19, 2 ≦ ν ≦ 5, and νβ <0.8), and the main phase is an intermediate crystal between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure A rare earth permanent magnet characterized by being an R′-LRE—Fe—Co based ferromagnetic compound having a structure.

In a more preferred specific example, a range of 0.1 ≦ νβ ≦ 0.5 is defined. Further, the density is preferably larger than when ν is 0. In a specific embodiment, the intermediate crystal structure with the TbCu ₇ crystal structure and ThMn ₁₂ type crystal structure, the TbCu ₇ crystal Fe atom pair of the rare earth element and dumbbell type was completely irregularly substituted This is an intermediate crystal structure between the structure and a ThMn ₁₂ type crystal structure in which rare earth elements and dumbbell type Fe atom pairs are regularly substituted. Such a structure can be characterized by the intensity of the superlattice diffraction peak of the XRD measurement. In a specific example, the intensity of the superlattice diffraction peak of the XRD measurement is characterized by an intermediate intensity between the superlattice diffraction peak intensities of the TbCu ₇ type crystal structure and the ThMn ₁₂ type crystal structure.

In a specific example of an embodiment, the rare earth permanent magnet is represented by a space group Immm, and the diffraction peak intensities of (310) and (002) in particular have a finite value in the space group Immm.

In a specific example of an embodiment, the R′-LRE—Fe—Co based ferromagnetic alloy described above includes a phase having a composition rich in rare earth elements and Cu, and the ratio of the rare earth elements constituting the phase is The atomic ratio is LRE> R ′.

Another aspect of the present invention provides a process A for preparing a molten alloy containing R ′, LRE, Fe, and Co, and cooling and solidifying the molten alloy so that at least a site occupied by rare earth elements in the alloy is obtained. Forming an R′-LRE-Fe—Co based ferromagnetic alloy including an R′—LRE—Fe—Co based ferromagnetic compound, which is a ferromagnetic compound partially substituted by Fe atom pairs; Step C for preparing a compound having a liquid phase composition, Step D for pulverizing an R′-LRE-Fe—Co based ferromagnetic alloy and a compound having a liquid phase composition, and pulverized R′-LRE-Fe—Co based ferromagnetic A step E of mixing the alloy and a compound having a liquid phase composition and a step F of densifying the magnetic powder of the R′-LRE—Fe—Co based ferromagnetic alloy in a state where the liquid phase is formed are included. As a result of these steps, the formula R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (where 0 <x <0.5, 0 <y <0.5, and 10 < z <19 and 0.01 ≦ α <0.5), or the formula R ′ _1-x LRE _{x + β} (Fe _1-y Co _y ) _z Cu _νβ (0 <x <0.5 and 0 <Y <0.5, 10 <z <19, 2 ≦ ν ≦ 5, and νβ <0.8).

In addition, another aspect of the present invention provides a process G for preparing a molten alloy containing R ′, LRE, Fe, Co, and Cu, and cooling and solidifying the molten alloy so that the rare earth element of the alloy is solidified. An R′-LRE-Fe—Co based ferromagnetic alloy including an R′—LRE—Fe—Co based ferromagnetic compound, which is a ferromagnetic compound in which at least a part of occupied sites is randomly substituted by Fe atom pairs, is formed. Step H, Step I for crushing the R′-LRE-Fe—Co based ferromagnetic alloy, and Step J for densifying the magnetic powder of the R′-LRE—Fe—Co based ferromagnetic alloy in a state where a liquid phase is formed including. As a result of these steps, the formula R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (where 0 <x <0.5, 0 <y <0.5, and 10 < z <19 and 0.01 ≦ α <0.5), or the formula R ′ _1-x LRE _{x + β} (Fe _1-y Co _y ) _z Cu _νβ (0 <x <0.5 and 0 <Y <0.5, 10 <z <19, 2 ≦ ν ≦ 5, and νβ <0.8).

More specifically, a heat treatment step K for heating the R′-LRE-Fe—Co based ferromagnetic alloy at 850 ° C. or lower is further included.

As another more preferable specific example, the step F or J is characterized by press molding at a temperature of 900 ° C. or lower.

According to the present invention, it is possible to provide a new ferromagnetic alloy and a manufacturing method thereof capable of solving the problem that occurs when the magnet of the ThMn ₁₂ type intermetallic compound generated by the rapid quenching method.

FIG. 3 is a structural diagram schematically showing a typical example of the crystal structure of an R ′ -LRE-Fe—Co based ferromagnetic compound in the present invention. FIG. 6 is a correspondence diagram showing the correspondence between the crystal structure of the R′—LRE—Fe—Co based ferromagnetic compound, the ThMn ₁₂ type crystal structure, and the TbCu ₇ type crystal structure in the present invention. It is a structural view showing a ThMn ₁₂ type crystal structure. It is a structural view showing a TbCu ₇ type crystal structure. FIG. 3 is a structural diagram showing an example of a crystal structure of an R′-LRE-Fe—Co based ferromagnetic compound in an example of the present invention. The flowchart which shows the manufacturing process of the manufacturing method of the rare earth permanent magnet of the Example of this invention. The flowchart which shows the manufacturing process of the manufacturing method of the rare earth permanent magnet of the Example of this invention. In an exemplary embodiment of the present _invention, it is a table showing the density of the molded body in _{Sm 0.45 Y 0.55 (Fe 0.83 Co} 0.17) 11 Ti 0.2 Cu x composition. In an exemplary embodiment of the present _invention, it is a table showing the density of the molded body in _{Sm 0.5 + x Y 0.5 (Fe} 0.83 Co 0.17) 11 Ti 0.2 Cu 2x composition. In an exemplary embodiment of the present _invention, it is a table showing the density of the molded body in _{Sm 0.5 + x Y 0.5 (Fe} 0.83 Co 0.17) 11 Ti 0.2 Cu 5x composition.

Hereinafter, embodiments will be described in detail with reference to the drawings. However, the present invention is not construed as being limited to the description of the embodiments below. Those skilled in the art will readily understand that the specific configuration can be changed without departing from the spirit or the spirit of the present invention.

In the structure of the invention described below, the same portions or portions having similar functions are denoted by the same reference numerals in different drawings, and redundant description may be omitted.

The position, size, shape, range, etc. of each component shown in the drawings and the like may not represent the actual position, size, shape, range, etc. in order to facilitate understanding of the invention. For this reason, the present invention is not necessarily limited to the position, size, shape, range, and the like disclosed in the drawings and the like.

[1. Composition and structure of R′-LRE-Fe—Co ferromagnetic compound]
The R′—LRE—Fe—Co based ferromagnetic alloy according to the present invention is an R′—LRE—Fe—Co based ferromagnetic alloy containing an R′—LRE—Fe—Co based ferromagnetic compound having a space group of Immm. . In this specification, “R ′” is a rare earth element containing at least Y (yttrium) or Gd (gadolinium). “LRE” is at least one rare earth element selected from La, Ce, Nd, Pr, and Sm. From the viewpoint of magnetic property values, Sm is particularly suitable for LRE. In this R'-LRE-Fe-Co based ferromagnetic compound, at least a part of occupied sites (possible occupying sites) of rare earth elements in the body-centered tetragonal ThMn ₁₂ type crystal structure is formed by a pair of Fe atoms (Fe dumbbells). It is a randomly substituted ferromagnetic compound. In other words, this R′-LRE—Fe—Co based ferromagnetic compound is constituted by an intermediate crystal structure between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure.

FIG. 1 schematically shows the crystal structure of an example of the R′-LRE—Fe—Co based ferromagnetic compound according to this example. In FIG. 1, the sites that can be occupied by the rare earth elements R ′, LRE, and Fe dumbbells are described with large circles overlapping with Fe dumbbells. More specifically, 2a site (gray circle) and 2d site (white circle) are shown as occupied sites of the rare earth elements R ′ and LRE. On the other hand, 4 g ₁ site (thick circled circle) and 4 g ₂ site (thin hatched circle) are shown as the sites occupied by the Fe dumbbell.

In the R'-LRE-Fe-Co ferromagnetic compound according to the present example, the Fe dumbbell can occupy the occupied sites of the rare earth elements R 'and LRE to some extent at random. That is, in the crystal structure of the R′-LRE—Fe—Co based ferromagnetic compound in this example, the Fe dumbbell is not completely randomly substituted with the rare earth element R ′.

The crystal structure in which the Fe dumbbell pair is completely randomly substituted with the rare earth element R ′ is a TbCu _7- type crystal structure. Therefore, in the X-ray diffraction pattern of the R′-LRE-Fe—Co based ferromagnetic compound according to this example, superlattice diffraction indicating the development of regularity from the TbCu ₇ type crystal structure to the ThMn ₁₂ type crystal structure is observed. Is done. However, the intensity of these superlattice diffraction peaks is weak compared to the intensity of the superlattice diffraction peaks generated from the ThMn ₁₂ type crystal structure without substitution of rare earth elements and Fe dumbbells. In particular, the diffraction peaks of (310) and (002) are suitable as indices in that they do not overlap with the intensity or other peaks in powder X-ray diffraction. These diffraction peaks cannot be observed in the TbCu ₇ type crystal structure. In addition, in the R′-LRE—Fe—Co based ferromagnetic compound according to the present example, a diffraction peak intensity weaker than the diffraction peak intensity observed in the ThMn ₁₂ type crystal structure is observed.

FIG. 2 shows that the crystal structure of the R′-LRE-Fe—Co ferromagnetic compound according to this example is an intermediate structure between the ThMn ₁₂ type crystal structure and the TbCu ₇ type crystal structure. Shown in relationship. The R′-LRE—Fe—Co based ferromagnetic compound according to this example continuously forms an intermediate structure between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure under the heat treatment conditions. Spatial County Immm is used to indicate a simple structure. By eliminating the 6-fold rotational symmetry around the c-axis of the TbCu ₇ type and the 4-fold rotational symmetry around the c-axis of the ThMn ₁₂ type, leaving the symmetry of the body center, this intermediate crystal structure can be transformed into a rare earth element. Can be expressed as a continuous replacement of Fe and dumbbell.

FIG. 3 corresponds to FIG. 2 and shows the relationship between the crystal structure of the R′-LRE-Fe—Co based ferromagnetic compound, ThMn ₁₂ type crystal structure, and TbCu ₇ type crystal structure according to this example. It is shown schematically for clarity.

FIG. 3A shows a ThMn ₁₂ type crystal structure. The crystal structure of the ThMn ₁₂ type is the space group I4 / mmm, and the crystal lattice constant is defined by a _tetra and c _tetra . In the ThMn ₁₂ type crystal structure, the Fe dumbbell is located on the Fe dumbbell line 301 in the occupied site of the rare earth element R, but not on the rare earth element line 302. Due to this regularity, peaks are observed in the ThMn ₁₂ type crystal structure.

FIG. 3B shows a TbCu ₇ type crystal structure. The crystal structure of TbCu ₇ type is the space group P6 / mmm, and the crystal lattice constant is defined by a _hex and c _hex in the figure. In the TbCu _7- type crystal structure, the Fe dumbbell can exist at an arbitrary position of the site occupied by the rare earth element R. That is, in the TbCu _7- type crystal structure, the occupation probability of the Fe dumbbell is not different between the Fe dumbbell line 301 and the rare earth element line 302. Therefore, as described above, no diffraction peak can be observed in the TbCu _7- type crystal structure.

FIG. 3C shows the crystal structure of the R′-LRE—Fe—Co based ferromagnetic compound according to this example. FIG. 3C is an intermediate structure between FIG. 3A and FIG. 3B, and the crystal structure of the TbCu ₇ type and the ThMn ₁₂ type can be expressed as a continuous change by the difference in substitution amount between the rare earth element and the Fe dumbbell pair. In addition, the symmetry of the space group Immm is introduced. The lattice constant of the space group Immm is defined as a _ortho , b _ortho , c _ortho . That is, by changing the parameters of the space group Immm (replacement amount of rare earth elements and Fe dumbbell pairs, lattice constant, site occupancy and internal coordinates), there is a TbCu _7- type crystal structure at one end, It is shown that there is a ThMn ₁₂ type crystal structure as the end structure. For this reason, the correspondence of the sites of the crystal lattice is as shown in FIG.

In the crystal structure of the R′-LRE—Fe—Co based ferromagnetic compound according to this example, the occupation probability of the Fe dumbbell is not equal between the Fe dumbbell line 310 and the rare earth element line 302. A crystal structure having such irregularity at the position of the Fe dumbbell and satisfying a _ortho = b _ortho in the lattice constant is referred to as “irregular ThMn ₁₂ type” in this specification. The irregular ThMn ₁₂ type crystal structure has fourfold symmetry around the c-axis, and the space group is I4 / mmm. In orthorhombic crystals, there is a prohibition of a _ortho ≠ b _ortho , but by removing this prohibition, a continuous change in crystal structure is expressed. In the structure of FIG. 3C, a weaker peak than that of the ThMn ₁₂ type is observed, but a stronger peak is observed than the TbCu ₇ type in which no substantial peak can be observed. That is, it can be said that this R′-LRE—Fe—Co based ferromagnetic compound has an intermediate structure between the ThMn ₁₂ type and the TbCu ₇ type.

The R′—LRE—Fe—Co based ferromagnetic compound according to this example has a composition expressed as R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α , where 0 <x <0. It is desirable that the composition ranges 5 and 0 <y <0.5 and 10 <z <19. From the viewpoint of increasing the magnetic anisotropy energy, the LRE is preferably Sm, and the magnetic anisotropy energy increases according to the amount of substitution. Therefore, it is desirable to replace LRE as much as possible, but when the amount of substitution of LRE is too large, the main phase is not produced in a sufficient amount for practical use. Moreover, partial substitution of Co is preferable from the viewpoint of improving magnetization at room temperature and improving magnetic anisotropy associated with an increase in Curie temperature. However, if the amount of substitution is too large, it is undesirable because it results in a decrease in magnetization and a decrease in magnetic anisotropy. Finally, it is desirable that the ratio between the rare earth element and the transition metal be generated in an amount sufficient for the main phase to be practically used. From the viewpoint of magnetic properties, a composition range of 0 ≦ x ≦ 0.5, 0.1 ≦ y ≦ 0.3, and 10.5 <z <14.0 is more desirable.

The R′-LRE-Fe—Co based ferromagnetic compound according to this example has a volume at room temperature when the composition is Y _0.6 Sm _0.4 (Fe _0.83 Co _0.17 ) _11.5 , for example. Since the magnetization has a value around 1.6T, the magnetic anisotropy field around 7T, and the Curie temperature around 520 ° C., it has a possibility as a hard magnetic phase. However, since it is a non-equilibrium phase, remarkable decomposition occurs by heat treatment at 900 ° C. or higher, and most of the main phase decomposes at 1000 ° C. or higher depending on the composition. Therefore, in order to put it into practical use as a bulk magnet, a method of densifying at 900 ° C. or lower is necessary. Desirably, a mechanism for aligning the crystal orientation in one direction is also required.

Since the present inventors search for an additive element that is difficult to dissolve in the main phase and forms a liquid phase at a high temperature, an element having a low affinity with the Fe element and a high affinity with the rare earth element is selected from the binary phase diagram. Selected. As a result, it was found that a liquid phase having a melting point of 900 ° C. or lower was generated by adding copper (Cu). It was confirmed that the coercive force may be improved as well as densification by adjusting the amount added. According to the inference from the binary phase diagram, it can be expected that silver (Ag) and bismuth (Bi) show such an effect.

Hereinafter, the liquid phase composition produced by the addition of Cu will be described.

[2. Liquid phase composition]
The sample in which the liquid phase is generated is rapidly solidified, and the tissue is subjected to energy dispersive X-ray spectroscopy (EDX) and powder X-ray diffraction (X-ray) using a scanning electron microscope (SEM). Analysis by diffraction (XRD) confirmed that a rare earth element, a Cu-rich Laves phase, and a CaCu _5- type near structure were generated. However, the “CaCu ₅ type neighborhood structure” referred to here has a CaCu ₅ type diffraction pattern, but is used to clearly indicate that the relative intensity ratios of the peaks are different, so that they are strictly different. That is, the composition of the liquid phase produced by adding Cu is RCu _γ (2 ≦ γ ≦ 5) X. In addition, the atomic ratio between R ′ and LRE constituting the rare earth element R is always LRE> R ′ in both the Laves phase and the CaCu ₅ type vicinity structure. It can be said that the light rare earth element LRE is easier to form a compound with the Cu element than the heavy rare earth element or the R ′ element having the characteristics thereof. Reversible endothermic peaks corresponding to the melting point of the liquid phase are present around 820 ° C. and 850 ° C. from DSC measurement of the temperature rise. Since these endothermic peaks change the heat amount according to the Cu addition amount but do not change the temperature, it can be said that the composition of the generated liquid phase does not change according to the Cu addition amount in the studied range.

When the casting composition is represented by R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (0 <x <0.5, 0 <y <0.5, 10 <z <19), α When <0.01, the amount of Cu added is too small to ensure a sufficient liquid phase necessary for densification. When α> 0.5, the amount of Cu added is too large, so the main phase is significantly decomposed. More desirably, the composition range is 0.01 ≦ α ≦ 0.5.

In consideration of the generation of the liquid phase, it is desirable to weigh the amount of rare earth elements larger than the desired composition of the R′-LRE—Fe—Co based ferromagnetic compound. It is particularly desirable to weigh more LRE. In that case, the casting composition is R ′ _1−x LRE _{x + β} (Fe _1−y Co _y ) _z Cu _νβ (0 <x < _0.5, 0 <y <0.5, 10 <z <19) When expressed, 2 <ν <5 is desirable. When ν ≦ 2, the amount of LRE introduced is large, and a phase of Th ₂ Ni ₁₇ type crystal structure or Th ₂ Zn ₁₇ type crystal structure is more easily generated than R′-LRE-Fe—Co based ferromagnetic compounds. When ≦ 5, since the amount of LRE introduced is small, the effect of increasing the amount of liquid phase generated cannot be sufficiently obtained, and as a result, an R′-LRE—Fe—Co based ferromagnetic compound having a desired composition cannot be obtained. A range of 3 ≦ ν ≦ 4 is more desirable. Furthermore, since the main phase ratio decreases as the liquid phase amount increases, it is desirable to adjust β according to the magnitude of ν so that νβ is <0.8.

Hereinafter, an example of an embodiment of a method for producing an R′-LRE—Fe—Co ferromagnetic alloy of the present invention will be described step by step. First, a manufacturing method in which an R′-LRE—Fe—Co based ferromagnetic alloy and a liquid phase alloy are separately prepared and densified in a mixed state will be described.

[3. Method for producing R'-LRE-Fe-Co based ferromagnetic alloy]
(A) Process for producing R'-LRE-Fe-Co master alloy R ', LRE, Fe and Co, or an alloy composed of two or more of them are mixed and dissolved in a vacuum or an inert gas. To produce a master alloy (melting casting method). The alloy composition is made uniform by melting. By using an R′-LRE—Fe—Co alloy having a known composition prepared in advance, there is an advantage that the composition can be easily adjusted during metal melting in the rapid solidification method. The composition deviation in the ingot of the produced R′—LRE—Fe—Co master alloy can be corrected in the step (B) described later. As another method, a plurality of alloys having different compositions can be separately produced and mixed in the step (B) described later.

The composition analysis of the R'-LRE-Fe-Co master alloy ingot can be performed by, for example, an inductively coupled plasma optical emission spectroscopy (ICP-OES) method. The compositional deviation can be suppressed by shortening the temperature raising time for dissolution or by adding a rare earth metal lump. In particular, when Sm is selected as the LRE, the latter is effective because the vapor pressure of Sm is high and it is easy to evaporate.

Instead of the above method, a reduction diffusion method in which an oxide or metal of a constituent element is mixed with granular calcium metal and heated in an inert gas atmosphere may be used. Since it does not involve a peritectic reaction, the formation of an Fe—Co phase that is soft magnetic can be suppressed, which is advantageous.

(B) Step of rapidly solidifying the mother alloy In this embodiment, the R′-LRE—Fe—Co mother alloy produced above is rapidly solidified to produce a rapidly solidified alloy. Examples of the rapid solidification method include a roll rapid cooling method such as a gas atomizing method, a single roll rapid cooling method, a twin roll rapid cooling method, a strip casting method, and a melt spinning method. Since the rare earth iron alloy is easily oxidized, it is preferable to rapidly cool it in a vacuum or in an inert atmosphere at a high temperature.

The disordered Th ₂ Ni ₁₇ type compound phase (R ′, LRE) ₂ (Fe, Co) ₁₇ has higher thermal stability than the R′-LRE-Fe—Co based ferromagnetic compound in this example. Even when the heat treatment step (K) described later is performed, the R'-LRE-Fe-Co based ferromagnetic compound in this example is not changed, and the irregular (R ', LRE) ₂ (Fe, Co) ₁₇ remains as it is. is there. Therefore, the generation of irregular (R ′, LRE) ₂ (Fe, Co) ₁₇ is suppressed at the time of rapid solidification in terms of securing the amount of R′—Fe—Co based ferromagnetic compound generated in this example. Is preferred. This is possible by increasing the cooling rate.

In the case where the melt spinning method using an air-cooled Cu single roll is used, in some embodiments, the roll peripheral speed is preferably set to 15 m / s or more. When the roll peripheral speed is 20 m / s or more, the R′-LRE—Fe—Co based ferromagnetic compound is produced at a rate of 50 wt% or more. By increasing the roll peripheral speed, the generation of an irregular Th ₂ Ni ₁₇ type compound phase can be suppressed, and the amount of R′—Fe—Co based ferromagnetic compound generated in this example increases. Therefore, it is more preferable to set the roll peripheral speed to 30 m / s or more.

On the other hand, depending on the heat treatment temperature in the heat treatment step (K) described later, the structure of the R′-LRE-Fe—Co based ferromagnetic compound in this example changes and thermal decomposition occurs. Therefore, depending on the heat treatment temperature in step (K), even if the roll peripheral speed is increased, the amount of R′-LRE—Fe—Co based ferromagnetic compound produced in this example does not change. From the viewpoint of productivity, the roll peripheral speed is preferably set to 50 m / s or less. Other embodiments of the invention are possible by non-equilibrium processes that produce metastable phases other than rapid solidification. For example, a nanoparticle process or a thin film process. Examples include a molecular beam epitaxy method, a sputtering method, an EB vapor deposition method, a reactive vapor deposition method, a laser ablation method, a vapor phase method such as a resistance heating vapor deposition method, a liquid phase method such as a microwave heating method, and a mechanical alloy method.

(C) Process for producing a sample having a liquid phase composition The production process basically conforms to the process (A). That is, R ′, LRE, Cu, or an alloy composed of two or more of them is mixed and melted in a vacuum or an inert gas to produce a master alloy (melting casting method). The alloy composition is made uniform by melting. In this case, it is important to adjust the composition so that LRE> R ′. As described above, the compositional deviation can be suppressed by shortening the heating time for melting or by adding a rare earth metal lump. In particular, when Sm is selected as the LRE, the latter is effective because the vapor pressure of Sm is high and it is easy to evaporate. As another embodiment of the present invention, melt sping may be used, and there is an advantage that a sample having a uniform liquid phase composition can be easily produced.

(D) Pulverization / Classification Step When densifying, an appropriate amount of liquid phase needs to be in contact with the R′-LRE-Fe—Co based ferromagnetic alloy. For this purpose, the samples prepared in the steps (B) and (C) must be pulverized to a certain particle size or less before the densification step (F) described later. The R′-LRE—Fe—Co ferromagnetic alloy prepared in the step (B) can be sufficiently pulverized to pulverize to 150 μm or less, and the liquid phase composition alloy prepared in the step (C) can be obtained as an R′— It is desirable to grind to a particle size smaller than that of the LRE-Fe-Co ferromagnetic alloy. By doing so, it is possible to sufficiently spread the liquid phase around the R′-LRE—Fe—Co based ferromagnetic alloy. On the other hand, if the particle size is too fine, problems such as oxidation and mold galling in the densification step (K) described later are not preferable. It is desirable to classify to a particle size of 1 μm to 150 μm. In order to suppress oxidation, it is preferable to work in an inert atmosphere in the glove box.

(E) Mixing step It is important to sufficiently mix the samples prepared in the step (B) and the step (C) in order to make the sample uniform in the liquid phase. Further, the liquid phase amount can be adjusted by this mixing step (E). A desirable liquid phase addition amount is 10 wt% or less based on the weight of the R′-LRE—Fe—Co based ferromagnetic alloy. A liquid phase of 10 wt% or more causes a decrease in the density of the main phase, which is not preferable in terms of magnetic characteristics, and has a problem that it is difficult to take out a sample from a mold. In addition, when the addition amount is 1 wt% or less, a liquid phase amount sufficient for densification cannot be secured.

(F) High-temperature densification step The densification step is desirably performed by applying pressure in a temperature range in which a liquid phase exists. The operating temperature for densification is preferably in the temperature range of 800 ° C to 900 ° C. When the temperature is 900 ° C. or higher, the mechanical strength of the mold is lowered, and the mold is deformed by a plurality of examinations. When the temperature exceeds 850 ° C., the R′-LRE-Fe—Co ferromagnetic compound begins to decompose into an irregular Th ₂ Ni ₁₇ type crystal structure and an Fe—Co phase, which are undesirable in terms of magnetic properties. become. On the other hand, the liquid phase has a melting point near 820 ° C. Therefore, it is more desirable to densify in the temperature range of 820 ° C to 850 ° C.

As the above-mentioned densification step, it is appropriate to use a cemented carbide die containing Co. Further, although resistance heating may be used as a heat source, a high frequency with a high temperature rise rate is desirable in consideration of composition shift due to evaporation of rare earth elements and adhesion inside the mold. Further, the discharge plasma sintering method may be used because it can be densified in a short time.

The pressure to be applied may be any pressure as long as the magnetic powder is sufficiently densified. For example, when examined at 3.7 MPa, it can be sufficiently densified.

Next, a production method for producing an R′-LRE-Fe—Co based ferromagnetic alloy to which Cu is added and generating a liquid phase at a high temperature to be densified will be described. The advantage of this method is that a uniform liquid phase distribution can be obtained in advance, and the crystal orientation can be aligned in one direction by performing die-up setting.

(G) Process for producing R'-LRE-Fe-Co-Cu master alloy R ', LRE, Fe, Co, Cu, or an alloy composed of two or more of them is mixed to form a vacuum or an inert gas Dissolve in to make a master alloy. The alloy composition is made uniform by melting. By using an R′-LRE—Fe—Co—Cu alloy having a known composition prepared in advance, there is an advantage that the composition can be easily adjusted when the metal is melted in the rapid solidification method. The composition deviation in the ingot of the produced R′—LRE—Fe—Co—Cu master alloy can be corrected in the step (H) described later. As another method, a plurality of alloys having different compositions can be separately produced and mixed in the step (H) described later.

The composition analysis of the R'-LRE-Fe-Co-Cu master alloy ingot can be performed, for example, by inductively coupled plasma emission spectroscopy. The compositional deviation can be suppressed by shortening the temperature raising time for dissolution or by adding a rare earth metal lump. In particular, when Sm is selected as the LRE, the latter is effective because the vapor pressure of Sm is high and it is easy to evaporate.

Instead of the above method, a reduction diffusion method in which an oxide or metal of a constituent element is mixed with granular calcium metal and heated in an inert gas atmosphere may be used. Since it does not involve a peritectic reaction, the formation of an Fe—Co phase that is soft magnetic can be suppressed, which is advantageous.

(H) Step of rapidly solidifying the mother alloy In this embodiment, the R′-LRE—Fe—Co—Cu mother alloy produced above is rapidly solidified to produce a rapidly solidified alloy. Examples of the rapid solidification method include a roll rapid cooling method such as a gas atomizing method, a single roll rapid cooling method, a twin roll rapid cooling method, a strip casting method, and a melt spinning method. Since the rare earth iron alloy is easily oxidized, it is preferable to rapidly cool it in a vacuum or in an inert atmosphere at a high temperature.

The disordered Th ₂ Ni ₁₇ type compound phase (R ′, LRE) ₂ (Fe, Co) ₁₇ has higher thermal stability than the R′-LRE-Fe—Co based ferromagnetic compound in this example. Even when the heat treatment step (K) described later is performed, the R'-LRE-Fe-Co based ferromagnetic compound in this example is not changed, and the irregular (R ', LRE) ₂ (Fe, Co) ₁₇ remains as it is. is there. Therefore, the generation of irregular (R ′, LRE) ₂ (Fe, Co) ₁₇ is suppressed at the time of rapid solidification in terms of securing the amount of R′—Fe—Co based ferromagnetic compound generated in this example. Is preferred. This is possible by increasing the cooling rate.

In the case where the melt spinning method using an air-cooled Cu single roll is used, in some embodiments, the roll peripheral speed is preferably set to 15 m / s or more. When the roll peripheral speed is 20 m / s or more, the R′-LRE—Fe—Co based ferromagnetic compound is produced at a rate of 50 wt% or more. By increasing the roll peripheral speed, the generation of an irregular Th ₂ Ni ₁₇ type compound phase can be suppressed, and the amount of R′—Fe—Co based ferromagnetic compound generated in this example increases. Therefore, it is more preferable to set the roll peripheral speed to 30 m / s or more.

On the other hand, depending on the heat treatment temperature in the heat treatment step (K) described later, the structure of the R′-LRE-Fe—Co based ferromagnetic compound in this example changes and thermal decomposition occurs. Therefore, depending on the heat treatment temperature in step (K), even if the roll peripheral speed is increased, the amount of R′-LRE—Fe—Co based ferromagnetic compound produced in this example does not change. From the viewpoint of productivity, the roll peripheral speed is preferably set to 50 m / s or less. Other embodiments of the invention are possible by non-equilibrium processes that produce metastable phases other than rapid solidification. For example, a nanoparticle process or a thin film process. Examples include a molecular beam epitaxy method, a sputtering method, an EB vapor deposition method, a reactive vapor deposition method, a laser ablation method, a vapor phase method such as a resistance heating vapor deposition method, a liquid phase method such as a microwave heating method, and a mechanical alloy method.

(I) Pulverization / Classification Step The particle diameter of the R′-LRE—Fe—Co ferromagnetic alloy at the time of densification is not limited as in the step (D) described above. It is sufficient if the R′—LRE—Fe—Co based ferromagnetic alloy enters the mold and the magnetic particles are sufficiently in contact with each other. For example, it is 500 μm or less. On the other hand, when performing die-up setting, it may depend on the amount of liquid phase to be produced, but it is appropriate to pulverize and classify to about 150 μm.

(J) High-temperature densification process As in the step (F) described above, the densification step is preferably performed by applying pressure in the temperature range where the liquid phase is in a liquid state. The operating temperature for densification is preferably in the temperature range of 800 ° C to 900 ° C. When the temperature is 900 ° C. or higher, the mechanical strength of the mold is lowered, and the mold is deformed by a plurality of examinations. When the temperature exceeds 850 ° C., the R′-LRE-Fe—Co ferromagnetic compound begins to decompose into an irregular Th ₂ Ni ₁₇ type crystal structure and an Fe—Co phase, which are undesirable in terms of magnetic properties. become. On the other hand, the liquid phase has a melting point near 820 ° C. Therefore, it is more desirable to densify in the temperature range of 820 ° C to 850 ° C.

As the above-mentioned densification step, it is appropriate to use a cemented carbide die containing Co. Further, although resistance heating may be used as a heat source, a high frequency with a high temperature rise rate is desirable in consideration of composition shift due to evaporation of rare earth elements and adhesion inside the mold. Further, the discharge plasma sintering method may be used because it can be densified in a short time.

The pressure to be applied may be any pressure as long as the magnetic powder is sufficiently densified. For example, when examined at 2.9 MPa, it can be sufficiently densified.

Furthermore, there is a case where the crystal orientation can be aligned in one direction by applying pressure to the compacted compact with a large-diameter mold so that the deformation rate is 50% or more. The operating temperature is preferably a temperature range in which a liquid phase is generated and decomposition of the main phase is not significant, and a temperature range of 820 ° C. to 850 ° C. is desirable.

(K) Heat treatment step The R′-LRE-Fe—Co based ferromagnetic compound according to this example is obtained from a TbCu ₇ type crystal structure in which a rare earth element and a dumbbell type Fe atom pair are completely irregularly substituted. The crystal structure is continuously changed by heat-treating to a ThMn _12- type crystal structure in which dumbbell-type Fe atom pairs are regularly substituted. Therefore, the heat treatment temperature and the heat treatment time are important in the sense of controlling the crystal structure of the R′-LRE—Fe—Co based ferromagnetic compound. A large magnetic anisotropy energy can be obtained by ordering into a ThMn _12- type crystal structure. Therefore, in order to optimize the structure of the R′-LRE—Fe—Co based ferromagnetic alloy according to the present example or the R′-LRE—Fe—Co based ferromagnetic compound according to the present example formed by the above method. In a preferred embodiment, heat treatment is performed. Holding the sample in a high temperature environment for a long time may cause evaporation of rare earth elements and oxidation of the sample, and may reduce productivity. For this reason, it is desirable to carry out the heat treatment step at a temperature at which uniform heat treatment can be performed in a relatively short time. The temperature of the heat treatment can be set between 600 ° C. and 1000 ° C., for example. The heat treatment time can be set within a range of, for example, 0.01 hours or more and less than 10 hours. Considering the ordering of the R'-LRE-Fe-Co based ferromagnetic compound from the TbCu ₇ type crystal structure to the ThMn ₁₂ type crystal structure, a high temperature is preferable, but the decomposition of the R'-LRE-Fe-Co based ferromagnetic compound is A heat treatment temperature of 850 ° C. or lower is more desirable because it cannot be ignored.

This heat treatment step (K) can also be performed during the high-temperature densification step (F) or (J). By doing so, the number of processes can be reduced and there is a production advantage. Specifically, this can be achieved by holding at a desired temperature either before or after the pressure application operation in the high-temperature densification step (F) or (J). The holding time at that time depends on the temperature, but is generally less than 1 hour.

Examples of the present invention will be specifically described below, but the present invention is not limited to these examples.

FIG. 4 shows a manufacturing process of the manufacturing method of the rare earth permanent magnet of this example.

(Process A)
First, the total composition is 4.2Y-3.5Sm-76.6Fe-15.7Co (at%) (chemical formula Sm _0.45 Y _0.55 (Fe _0.83 Co _0.17 ) ₁₂ ) In order to obtain a raw material alloy having a weight of 1 kg, Y (purity 99.9%), Sm (purity 99.9%), electrolytic iron (purity 99.9%) and electrolytic cobalt (purity 99.9%) were weighed, respectively. . In consideration of evaporation of Y and Sm at high temperature, 63.5 g of Y, so that Y is 3 mass% and Sm is 5 mass% higher than the target composition 7.7Y-76.6Fe-15.7Co, 89.6 g Sm, 704.3 g Fe, and 148.7 g Co were weighed. The weighed metals were mixed, put into an alumina crucible, and melted by high frequency melting. Thereafter, the molten metal was spread on a water-cooled copper hearth and solidified to obtain an alloy ingot. As a result of analyzing the produced alloy ingot using an ICP analyzer (manufactured by Shimadzu Corporation: ICPV-1017), the composition was 3.8Y-3.4Sm-77.5Fe-15.3Co (at%). .

For the ingot having the composition of 3.8Y-3.4Sm-77.5Fe-15.3Co thus obtained, the overall composition is, for example, the chemical formula Sm _0.45 Y _0.55 (Fe _0.83 Co _0.17 ) when ₁₂ of, Y of the metal mass 0.063 g, was added weighed slug 0.016g of metal masses 0.030g and Co of Sm, quartz tapping with holes them in the bottom (0.8 mm) I put it in the tube. 3.8Y-3.4Sm-77.5Fe-15.3Co ingot, Y metal lump, Sm metal lump, and quartz hot metal pipe filled with Co metal lump are made into a high frequency induction heating type amorphous metal production furnace (Nisshin The product was introduced into Giken Co., Ltd., and the ingot and the metal lump were heated and dissolved by applying a high frequency electric field in an Ar atmosphere of 75 kPa.

Prepare an ingot near the desired composition in the same manner as described above, and heat the sample with the entire composition adjusted by adding an appropriate amount of the metal mass that is insufficient in Y, Sm, Fe, Co in the same procedure as above. And dissolved. The composition was adjusted in the range of Y _1-x Sm _x (Fe _1-y Co _y ) _z (0 <x <0.5, 0 <y <0.5, 10 <z <19) as a chemical formula. Hereinafter, in this example, the alloy composition is expressed by a chemical formula.

(Process B)
After confirming that the Y—Sm—Fe—Co alloy was sufficiently dissolved in step A, the molten metal was emitted onto a copper roll (roll diameter: 250 mm) rotating at high speed with Ar at a tapping pipe pressure of 100 kPa, and then rapidly solidified. A ribbon-like alloy (hereinafter referred to as an ultra-quenched ribbon) was produced. In this embodiment, a roll peripheral speed of 40 m / s is set as a basic condition. By increasing the roll peripheral speed, it is possible to suppress the generation of an irregular Th ₂ Ni ₁₇ type crystal structure and Fe—Co in an as-spun sample (sample not subjected to heat treatment after rapid solidification). In this specification, the cooling rate of the molten alloy is expressed by “roll peripheral speed”, but the cooling rate depends on the thermal conductivity, heat capacity, atmospheric pressure, tap pipe pressure, etc. of the roll used for cooling. Can also change. When a roll of a material or size different from the roll used in the examples of the present specification is used, the preferred range of the roll peripheral speed naturally varies.

(Process C)
In order to prepare a sample having a liquid phase composition, the melt spinning method was used in the same manner as in Step B described above. In order to prepare a sample in the composition range of Sm _0.7 Y _0.3 Cu _β (2 ≦ β ≦ 5), a quartz hot metal pipe into which a Y metal block, an Sm metal block and a Cu metal block are charged is a high frequency induction heating type. In an amorphous metal fabrication furnace (manufactured by Nisshin Giken Co., Ltd.), ingot and metal lump are heated and melted by applying a high-frequency electric field in an Ar atmosphere of 75 kPa, A molten metal was emitted onto a copper roll rotating at a high speed of 20 m / s and rapidly solidified to produce a ribbon-like alloy. For example, when a liquid phase sample having a composition of Sm _0.7 Y _0.3 Cu ₄ was prepared, 2.726 g of Sm metal lump, 0.691 g of Y metal lump, and 6.583 g of Cu metal lump were weighed.

(Process D)
The ultra-quenched ribbon produced in Step B and Step C was pulverized in a glove box under an Ar atmosphere using a pulverizer (Osaka Chemical Co., Ltd.). The Y—Sm—Fe—Co-based ultra-quenched ribbon obtained in step B was classified into magnetic particles having a particle diameter of 150 μm to 75 μm. In addition, the ultra-quenched ribbon with Sm _0.7 Y _0.3 Cu _β (2 ≦ β ≦ 5) composition obtained in Step C was pulverized to 20 μm or less.

(Process E)
Two Y-Sm-Fe-Co-based ultra-quenched ribbons obtained in step D and an ultra-quenched ribbon with a Sm _0.7 Y _0.3 Cu _β (2 ≦ β ≦ 5) composition were rotated into a V-shaped container. The mixture was introduced into a mixing machine and mixed uniformly (hereinafter, mixed magnetic powder). At that time, 5 wt% of the ultra-quenched ribbon with the composition of Sm _0.7 Y _0.3 Cu _β (2 ≦ β ≦ 5) was added with respect to the weight of the Y-Sm—Fe—Co-based ultra-quenched ribbon.

(Process F)
3 g of the mixed magnetic powder produced in Step E is put into a cemented carbide die (5φ) with a thermocouple welded, sandwiching a release carbon sheet, and a high-frequency induction heating type hot working device (Nisshin Giken) And heated by applying a high-frequency electric field in an Ar atmosphere of 75 kPa. After raising the temperature to 825 ° C. in 1 minute and holding for 15 minutes, a pressure of 2.9 MPa was applied for 3 minutes, and the pressure was released to cool.

It was confirmed from the SEM-EDX analysis that a Cu-rich phase having a higher Sm concentration than Y in terms of atomic ratio was formed on the formed body. Further, as a result of pulverizing the obtained molded body and performing XRD measurement (XRD: X-ray Diffraction (X-ray diffraction method)), diffraction peaks of (301) and (002) belonging to the main phase were observed. It was confirmed that the density of the obtained compact was higher in the magnetic powder packing density than the sample in which no liquid phase was generated.

(Process K)
After the ultra-quenched ribbon, powder or molded body produced after Step B, Step D or Step F is wrapped in Nb foil and loaded into a quartz tube having an Ar flow atmosphere, the quartz tube is set to a predetermined temperature in advance. It is also possible to put it in a tubular furnace and hold it. Thereafter, the quartz tube was dropped into water and sufficiently cooled. The heat treatment in the Ar flow can suppress evaporation of the rare earth element more than the heat treatment in the vacuum.

FIG. 5 shows a manufacturing process of the manufacturing method of the rare earth permanent magnet of the present embodiment.

(Process G)
The Y—Sm—Fe—Co-based alloy prepared in Step A of Example 1 was used. For an ingot having a composition of 3.8Y-3.4Sm-77.5Fe-15.3Co, the overall composition is, for example, the chemical formula Sm _0.45 Y _0.55 (Fe _0.83 Co _0.17 ) ₁₁ Cu _{0 In} the case of 0.2, 0.115 g of Y metal lump, 0.104 g of Sm metal lump, 0.015 g of Co metal lump and 0.167 g of Cu metal lump are weighed and added to the bottom (0 .8 mmφ) was put into an open quartz tap pipe. 3.8Y-3.4Sm-77.5Fe-15.3Co ingot, Y metal lump, Sm metal lump, Co metal lump and quartz metal hot metal tube charged with Cu metal lump are produced with high frequency induction heating type amorphous metal It was introduced into a furnace, and the ingot and the metal lump were heated and melted by applying a high frequency electric field in an Ar atmosphere of 75 kPa.

Prepare an ingot near the desired composition in the same manner as described above, and heat the sample with the entire composition adjusted by adding an appropriate amount of the metal mass that is insufficient in Y, Sm, Fe, Co in the same procedure as above. And dissolved. The composition is represented by the chemical formula Y _1-x Sm _x (Fe _1-y Co _y ) _z Cu _α (0 <x <0.5, 0 <y <0.5, 10 <z <19, 0 ≦ α ≦ 0. Adjustment was made within the range of 6). Hereinafter, in this example, the alloy composition is expressed by a chemical formula.

(Process H)
After confirming that the Y—Sm—Fe—Co—Cu alloy was sufficiently dissolved in Step G, the molten metal was ejected onto a copper roll (roll diameter: 250 mm) rotating at high speed with Ar having a tapping pressure of 100 kPa. A ribbon-like alloy (hereinafter referred to as an ultra-quenched ribbon) was produced by rapid solidification. In this embodiment, a roll peripheral speed of 40 m / s is set as a basic condition. By increasing the roll peripheral speed, it is possible to suppress the generation of an irregular Th ₂ Ni ₁₇ type crystal structure and Fe—Co in an as-spun sample (sample not subjected to heat treatment after rapid solidification).

(Process I)
The ultra-quenched ribbon produced in Step H was pulverized in a glove box under an Ar atmosphere using a pulverizer (Osaka Chemical Co., Ltd.). The Y—Sm—Fe—Co—Cu ultra-quenched ribbon obtained in Step H was classified into magnetic particles having a particle diameter of 150 μm to 75 μm.

(Process J)
3 g of Y-Sm-Fe-Co-Cu-based ultra-quenched ribbon is put into a Co-containing hard metal mold (5φ) with a carbon sheet for mold release and a thermocouple welded. The sample was introduced into a hot working apparatus and heated by applying a high frequency electric field in an Ar atmosphere of 75 kPa. After raising the temperature to 830 ° C. in 2 minutes and holding for 4 minutes, a pressure of 2.9 MPa was applied for 3 minutes, and the pressure was released to cool. Table 2 shows the density and characteristics of the prepared sample.

The molded body thus produced was similarly placed in a Co-containing carbide die (12φ) welded with a thermocouple, introduced into a high-frequency induction heating type hot working apparatus, heated to 830 ° C. in 1 minute, and 2.9 MPa. Was applied for 3 minutes, and the pressure was released to cool.

It was confirmed from the SEM-EDX analysis that a Cu-rich phase having a higher Sm concentration than Y in terms of atomic ratio was formed on the formed body. Further, as a result of pulverizing the obtained molded body and performing XRD measurement, diffraction peaks of (301) and (002) belonging to the main phase were observed. It was confirmed that the density of the obtained compact was higher in the magnetic powder packing density than the sample in which no liquid phase was generated.

(Process K)
The ultra-quenched ribbon or molded body produced after Step H or Step J is wrapped in Nb foil, loaded into a quartz tube having an Ar flow atmosphere, and then the quartz tube is placed in a tubular furnace set at a predetermined temperature in advance. The process of holding can also be performed. Thereafter, the quartz tube was dropped into water and sufficiently cooled. The heat treatment in the Ar flow can suppress evaporation of the rare earth element more than the heat treatment in the vacuum.

FIG. 6 is a table showing the density of the compact in the Sm _0.45 Y _0.55 (Fe _0.83 Co _0.17 ) ₁₁ Ti _0.2 Cu _x composition produced by the above process. As the amount of Cu introduced increases, the density tends to improve. The maximum density was shown at x to 0.3, and the density slightly decreased near x to 0.3. According to this result, it is possible to control the density by controlling the amount of Cu introduced in the range of 0.1 to 0.5, preferably 0.1 to 0.4, and more preferably 0.2 to 0.4. The effect of improvement can be obtained.

The process of the present embodiment is almost the same as the process of the second embodiment shown in FIG. 5, but different points will be particularly described.

(Process G)
The Y—Sm—Fe—Co-based alloy prepared in Step A of Example 1 was used. For an ingot having a composition of 3.8Y-3.4Sm-77.5Fe-15.3Co, the overall composition is, for example, the chemical formula Sm _0.55 Y _0.55 (Fe _0.83 Co _0.17 ) ₁₁ Cu _{0 In} the case of _.4 , 0.111 g of Y metal lump, 0.294 g of Sm metal lump, 0.015 g of Co metal lump and 0.327 g of Cu metal lump are weighed and added to the bottom (0 .8 mmφ) was put into an open quartz tap pipe. 3.8Y-3.4Sm-77.5Fe-15.3Co ingot, Y metal lump, Sm metal lump, Co metal lump and quartz metal hot metal tube charged with Cu metal lump are produced with high frequency induction heating type amorphous metal It was introduced into a furnace, and the ingot and the metal lump were heated and melted by applying a high frequency electric field in an Ar atmosphere of 75 kPa.

Prepare an ingot near the desired composition in the same manner as described above, and heat the sample with the entire composition adjusted by adding an appropriate amount of the metal mass that is insufficient in Y, Sm, Fe, Co in the same procedure as above. And dissolved. The composition is Y _1-x Sm _{x + β} (Fe _1-y Co _y ) _z Cu _νβ (0 <x < _0.5, 0 <y <0.5, 10 <z <19, 2 ≦ ν ≦ 5, Adjustment was made in the range of νβ ≦ 0.8). Hereinafter, in this example, the alloy composition is expressed by a chemical formula.

(Process H)
After confirming that the Y—Sm—Fe—Co—Cu alloy was sufficiently dissolved in Step G, the molten metal was ejected onto a copper roll (roll diameter: 250 mm) rotating at high speed with Ar having a tapping pressure of 100 kPa. A ribbon-like alloy (hereinafter referred to as an ultra-quenched ribbon) was produced by rapid solidification. In this embodiment, a roll peripheral speed of 40 m / s is set as a basic condition. By increasing the roll peripheral speed, it is possible to suppress the generation of an irregular Th ₂ Ni ₁₇ type crystal structure and Fe—Co in an as-spun sample (sample not subjected to heat treatment after rapid solidification).

(Process I)
The ultra-quenched ribbon produced in Step H was pulverized in a glove box under an Ar atmosphere using a pulverizer (Osaka Chemical Co., Ltd.). The Y—Sm—Fe—Co—Cu ultra-quenched ribbon obtained in Step H was classified into magnetic particles having a particle diameter of 150 μm to 75 μm.

(Process J)
3 g of Y-Sm-Fe-Co-Cu-based ultra-quenched ribbon is put into a Co-containing hard metal mold (5φ) with a carbon sheet for mold release and a thermocouple welded. The sample was introduced into a hot working apparatus and heated by applying a high frequency electric field in an Ar atmosphere of 75 kPa. The temperature was raised to 800 ° C. in 2 minutes and held for 4 minutes, and then a pressure of 2.9 MPa was applied for 3 minutes, and the pressure was released to cool. Table 2 shows the density and characteristics of the prepared sample.

The molded body thus produced was similarly placed in a Co-containing carbide die (12φ) welded with a thermocouple, introduced into a high-frequency induction heating type hot working apparatus, heated to 830 ° C. in 2 minutes, and 2.9 MPa. Was applied for 3 minutes, and the pressure was released to cool.
It was confirmed from the SEM-EDX analysis that a Cu-rich phase having a higher Sm concentration than Y in terms of atomic ratio was formed in the molded body thus prepared. In addition, as a result of pulverizing the obtained molded body and performing XRD measurement, diffraction peaks (301) and (002) belonging to the main phase (which may be X-ray diffraction, electron beam diffraction, or neutron diffraction) were observed. . It was confirmed that the density of the obtained compact was higher in the magnetic powder packing density than the sample in which no liquid phase was generated.

(Process K)
The ultra-quenched ribbon or molded body produced after Step H or Step J is wrapped in Nb foil, loaded into a quartz tube having an Ar flow atmosphere, and then the quartz tube is placed in a tubular furnace set at a predetermined temperature in advance. The process of holding can also be performed. Thereafter, the quartz tube was dropped into water and sufficiently cooled. The heat treatment in the Ar flow can suppress evaporation of the rare earth element more than the heat treatment in the vacuum.

FIG. 7 is a table showing the density of the compact in the composition of Sm _{0.5 + x} Y _0.5 (Fe _0.83 Co _0.17 ) ₁₁ Ti _0.2 Cu _2x produced by the above process. As the amount of Cu introduced increases, the density tends to improve. The maximum density was shown at x˜0.2.

Similarly, FIG. 8 shows the density of the compact in the composition of Sm _{0.5 + x} Y _0.5 (Fe _0.83 Co _0.17 ) ₁₁ Ti _0.2 Cu _5x . Similarly, the density tends to improve as the amount of Cu introduced increases. The maximum density was shown at x˜0.06.

From the above results, it was found that the density of the molded body can be increased by introducing Cu and adjusting the amount of introduction as compared with the case where Cu is not introduced.

According to this example, by adding a small amount of Cu to the quaternary system whose main component is Sm—Y—Fe—Co, a liquid phase having a melting point of about 820 ° C. can be generated, and a bulk magnet is manufactured. It was possible to promote densification. Moreover, since a liquid phase does not generate | occur | produce in the sample which has not introduce | transduced Cu, when observed with SEM, there are comparatively many gaps between magnetic powders. On the other hand, the sample introduced with Cu has a material structure with almost no gap due to the influence of liquid phase generation. According to this result, in this example, it is considered that the bulk magnet can be densified and the mechanical strength can be improved.

The present invention is not limited to the above-described embodiment, and includes various modifications. For example, a part of the configuration of one embodiment can be replaced with the configuration of another embodiment, and the configuration of another embodiment can be added to the configuration of one embodiment. Further, it is possible to add, delete, and replace the configurations of other embodiments with respect to a part of the configurations of the embodiments.

The R′-LRE—Fe—Co—Cu based ferromagnetic alloy of the present invention can be suitably used for a bulk magnet, for example.

301: Fe dumbbell line 302: Rare earth element line

Claims (15)

The composition of the R′-LRE—Fe—Co based ferromagnetic alloy (R ′ is at least one selected from Y and Gd, and LRE is at least one selected from La, Ce, Nd, Pr and Sm),
Formula R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (where 0 <x <0.5 and 0 <y <0.5 and 10 <z <19 and 0.01 ≦ α <0.5)
A rare earth permanent magnet, wherein the main phase is an R′-LRE—Fe—Co based ferromagnetic compound having an intermediate crystal structure between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure. Intermediate and the crystal structure, the TbCu ₇ crystal structure and rare earth element and dumbbell Fe atom pair of the rare earth element and dumbbell type was completely irregularly substitution of the of the TbCu ₇ crystal structure and ThMn ₁₂ type crystal structure Is an intermediate crystal structure with a regularly substituted FeMn pair of ThMn ₁₂ crystal structure, and the superlattice diffraction peak intensity of XRD measurement is the TbCu ₇ crystal structure and the ThMn ₁₂ crystal structure The rare earth permanent magnet according to claim 1, wherein the rare earth permanent magnet is characterized by being an intermediate intensity of the intensity of the superlattice diffraction peak. 3. The rare earth permanent magnet according to claim 2, wherein the rare earth permanent magnet is expressed by a space group Immm, and the diffraction peak intensities of (310) and (002) have a finite value in the space group Immm. The R′-LRE—Fe—Co based ferromagnetic alloy includes a phase having a composition rich in rare earth elements and Cu, and the ratio of the rare earth elements constituting the phase is LRE> R ′ in atomic ratio. The rare earth permanent magnet according to claim 1, characterized in that: The rare earth permanent magnet according to claim 1, wherein 0.1 ≦ α ≦ 0.4. 2. The rare earth permanent magnet according to claim 1, wherein a density of the molded body is larger than that when α is set to 0. The composition of the R′-LRE—Fe—Co based ferromagnetic alloy (R ′ is at least one selected from Y and Gd, and LRE is at least one selected from La, Ce, Nd, Pr and Sm),
Formula R ′ _1-x LRE _{x + β} (Fe _1-y Co _y ) _z Cu _νβ (0 <x <0.5 and 0 <y <0.5 and 10 <z <19 and 2 ≦ ν ≦ 5 and νβ <0.8)
A rare earth permanent magnet, wherein the main phase is an R′-LRE—Fe—Co based ferromagnetic compound having an intermediate crystal structure between a TbCu ₇ type crystal structure and a ThMn ₁₂ type crystal structure. Intermediate and the crystal structure, the TbCu ₇ crystal structure and rare earth element and dumbbell Fe atom pair of the rare earth element and dumbbell type was completely irregularly substitution of the of the TbCu ₇ crystal structure and ThMn ₁₂ type crystal structure Is an intermediate crystal structure with a regularly substituted FeMn pair of ThMn ₁₂ crystal structure, and the superlattice diffraction peak intensity of XRD measurement is the TbCu ₇ crystal structure and the ThMn ₁₂ crystal structure The rare earth permanent magnet according to claim 7, wherein the rare earth permanent magnet is characterized by being an intermediate intensity of the intensity of the superlattice diffraction peak. 9. The rare earth permanent magnet according to claim 8, wherein the rare earth permanent magnet is expressed by a space group Immm, and the diffraction peak intensities of (310) and (002) have a finite value in the space group Immm. The R′-LRE—Fe—Co based ferromagnetic alloy includes a phase having a composition rich in rare earth elements and Cu, and the ratio of the rare earth elements constituting the phase is LRE> R ′ in atomic ratio. The rare earth permanent magnet according to claim 7, wherein: The rare earth permanent magnet according to claim 7, wherein 0.1 ≦ νβ ≦ 0.5. The rare earth permanent magnet according to claim 7, wherein the density of the molded body is larger than that in the case where ν is 0. Preparing a melt of an alloy containing R ′, LRE, Fe and Co;
By cooling and solidifying the molten metal of the alloy, an R′-LRE-Fe—Co-based strong which is a ferromagnetic compound in which at least a part of the rare earth element occupation sites of the alloy is randomly substituted by Fe atom pairs. Forming an R′-LRE—Fe—Co based ferromagnetic alloy containing a magnetic compound;
Step C for producing a compound having a liquid phase composition;
Crushing the R′-LRE—Fe—Co based ferromagnetic alloy and the compound having the liquid phase composition;
Mixing the pulverized R′-LRE-Fe—Co ferromagnetic alloy with a compound having a liquid phase composition; and
Including a step F of densifying the magnetic powder of the R′-LRE—Fe—Co based ferromagnetic alloy in a state where a liquid phase is formed,
Or
Preparing a melt of an alloy containing R ′, LRE, Fe, Co and Cu;
By cooling and solidifying the molten metal of the alloy, an R′-LRE-Fe—Co-based strong which is a ferromagnetic compound in which at least a part of the rare earth element occupation sites of the alloy is randomly substituted by Fe atom pairs. Forming an R′—LRE—Fe—Co ferromagnetic alloy containing a magnetic compound;
Crushing R′-LRE—Fe—Co based ferromagnetic alloy;
Including the step J of densifying the magnetic powder of the R′-LRE—Fe—Co based ferromagnetic alloy in a state in which a liquid phase is formed,
Formula R ′ _1-x LRE _x (Fe _1-y Co _y ) _z Cu _α (where 0 <x <0.5 and 0 <y <0.5 and 10 <z <19 and 0.01 ≦ α <0.5), or the formula R ′ _1-x LRE _{x + β} (Fe _1-y Co _y ) _z Cu _νβ (0 <x <0.5 and 0 <y <0.5, and 10 <z <19, 2 ≦ ν ≦ 5, and νβ <0.8). The method for producing a rare earth permanent magnet according to claim 13, comprising a heat treatment step K in which the R'-LRE-Fe-Co based ferromagnetic alloy is heated at 850 ° C or lower. The method of manufacturing a rare earth permanent magnet according to claim 13, wherein the step F or J is press-molded at a temperature of 900 ° C or lower.

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