JP2020136333A - Sintered body for rare earth magnet and method for manufacturing the same - Google Patents

Abstract

To provide a sintered body for a rare earth magnet, which is less in the quantity of foreign phases having an adverse effect on a magnetic property, and a method for manufacturing the sintered body.SOLUTION: In an exemplary embodiment, a sintered body for a rare earth magnet disclosed herein has a whole composition given by the following compositional formula: R1R2(FeCo)TiCuO, where R1 is Y or Y and Gd, Y is 50 mol% or more to a whole quantity of R1, R2 is at least one kind selected from a group consisting of Sm, La, Ce, Nd and Pr, but it necessarily includes Sm, Sm is 50 mol% or more to a whole quantity of R2, and x, y, z, w, α and β satisfy 0.3≤x≤0.9, 0≤y≤0.4, 0.38≤z≤0.70, 7≤w≤12, 0≤α≤0.70, 0.02≤β≤0.5 and 0≤1-x-2z/3-0.092α-8β/15≤0.05. In the sintered body, a phase having a ThMntype crystalline structure makes a main phase.SELECTED DRAWING: Figure 2

Description

The present invention relates to a sintered body for rare earth magnets and a method for producing the same.

Permanent magnets are used in various motors such as automobile parts, industrial machines, and home appliances.

A typical high-performance magnet is an Nd-Fe-B magnet. Nd-Fe-B magnets are mainly used in drive motors of electric vehicles (EV, HV, PHV, etc.) and hybrid vehicles. Needs for further high efficiency and miniaturization of motors are increasing, and the development of permanent magnets having higher magnetic properties is expected.

One of the Nd-Fe-B system in the main phase alloy of the permanent magnets than the magnetic properties of the magnet candidate, RT ₁₂ compounds having ThMn ₁₂ type crystal structure or a similar structure has been noted. The RT ₁₂ compound is a compound that constitutes the main phase of an Nd-Fe-B magnet. R ₂ T ₁₄ B (R is at least one kind of rare earth element, T is one or more iron group transition metal elements containing at least Fe. ) High magnetic properties are expected because it contains a higher concentration of iron group transition metals. Hereinafter sometimes described as 1-12 phase phase consisting _{RT 12} compounds having ThMn ₁₂ type crystal structure or a similar structure.

Patent Document 1 discloses a rare earth permanent magnet in which a part of Fe, which is a T element, is partially replaced by Ti, which is a structural stabilizing element, to improve thermal stability in exchange for high magnetization. There is.

Patent Document 2 states that by partially substituting the R element of an RFe _12- based compound with a substitution element such as Zr or Hf, the amount of the substitution element such as Ti that replaces the transition metal element is reduced to achieve saturation magnetization. A rare earth permanent magnet in which the ThMn _12- type crystal structure is stabilized while maintaining the structure is disclosed.

Further, Patent Document 3 discloses an R'-Fe-Co-based ferromagnetic alloy in which Y or Gd is selected as a part of the R element of RFe ₁₂ , and the R'-Fe-Co-based ferromagnetic alloy is disclosed. It is described that the alloy has a ThMn _12- type crystal structure produced by an ultra-quenching method and thus exhibits high magnetic properties.

Further, Patent Document 4 describes that the addition of Cu produces a phase having a non-magnetic and low melting point of 1-4 composition (SmCu _4- phase), which enables sintering and high coercive force. ..

Further, in Patent Document 5, at least one of Sm ₅ Fe ₁₇ system phase, SmCo ₅ system phase, Sm ₂ O ₃ system phase, and Sm ₇ Cu ₃ system phase is used as a subphase with respect to the main phase of ThMn ₁₂ type. It is described that high coercive force can be increased by including it.

Further, Patent Document 6 describes that by adding Cu, a liquid phase is formed and a dense bulk body can be formed.

Further, in Patent Document 7, by producing a ferromagnetic alloy having a ThMn ₁₂ type phase containing Y as a main phase by a strip casting method, an alloy having less non-uniformity of the main phase composition and a high main phase ratio Is stated to be obtained.

Further, Patent Document 8 describes that a high saturation magnetization and an anisotropic magnetic field can be obtained with a magnet material having a ThMn ₁₂ type phase containing Y as a main phase.

Japanese Unexamined Patent Publication No. 64-76703 Japanese Unexamined Patent Publication No. 4-322406 Japanese Unexamined Patent Publication No. 2015-156436 Japanese Unexamined Patent Publication No. 2001-189206 JP-A-2017-112300 International Publication No. 2016/162990 Japanese Unexamined Patent Publication No. 2018-10321 JP-A-2018-125512

As one of the conditions for the sintered body used for high-performance magnets, it is necessary to have a structure having few different phases that adversely affect the magnetic properties. If a soft magnetic phase typified by the bcc-Fe phase is present in the sintered body, the soft magnetic phase becomes the starting point of the magnetization reversal and the magnetization reversal easily proceeds, so that coercive force, squareness, residual magnetic flux density, etc. The magnetic properties are significantly reduced. Therefore, there is a need for a sintered body in which such a soft magnetic heterogeneous phase does not exist as much as possible.

The rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to element substitution of Fe by Ti, but since the amount of Fe substitution by Ti is large, the magnetization is reduced by that amount, and sufficient magnetic properties are provided. I can't get it.

On the other hand, in the rare earth permanent magnet described in Patent Document 2, although the ThMn ₁₂ structure is stabilized by substituting the transition metal element with Ti or the like, the effect is not always sufficient.

Since the R'-Fe-Co-based ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structure stabilizing element M (Ti, etc.), it has high magnetization, large magnetic anisotropy, and high Curie temperature. Although it has been obtained, since it is a non-equilibrium phase, the main phase compound may decompose in a densification process at high temperature such as sintering.

In the rare earth magnet described in Patent Document 4, the magnetic property value may not be high because the amount of Ti added is large.

In the rare earth magnet described in Patent Document 5, when the rare earth rich subphase Sm ₇ Cu ₃ is used, the reaction between the main phase and Sm ₇ Cu ₃ during the heat treatment may generate a rare earth rich phase than the main phase. I am concerned.

In the rare earth magnet described in Patent Document 6, since the Fe element is not replaced with the structure stabilizing element M, high magnetization, large magnetic anisotropy and high Curie temperature can be obtained, and the density as a bulk compound is high. Since it is a non-equilibrium phase, the main phase compound may decompose in a process at a high temperature such as sintering at 1000 ° C. or higher.

In the composition of the ferromagnetic alloy described in Patent Document 7 and the magnet material described in Patent Document 8, oxygen is prioritized over rare earth elements because the influence of oxygen inevitably mixed in the sintered body manufacturing process is not taken into consideration. There is a concern that the main phase will be decomposed and a soft magnetic phase such as a bcc-Fe phase will be generated.

The embodiments of the present disclosure provide a sintered body for a rare earth magnet having few different phases that adversely affect the magnetic properties, and a method for producing the same.

In an exemplary embodiment, the overall composition of the sintered body for rare earth magnets having a phase having a ThMn _12- type crystal structure of the present disclosure as a main phase is represented by the following composition formula (1), and R1 _{1-x. R2 x (Fe 1-y Co} y) w-z Ti z Cu α O β (1), R1 is Y or Y and Gd, Y is at least 50 mol% of the total R1, R2 is Sm, La, It is at least one selected from the group consisting of Ce, Nd and Pr, always contains Sm, Sm is 50 mol% or more of the total R2, and x, y, z, w, α, and β are, respectively. 0.3 ≦ x ≦ 0.9, 0 ≦ y ≦ 0.4, 0.38 ≦ z ≦ 0.70, 7 ≦ w ≦ 12, 0 ≦ α ≦ 0.70, 0.02 ≦ β ≦ 0. 5 and 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.05 are satisfied.

In one embodiment, 0≤1-x-2z / 3-0.092α-8β / 15≤0.03
To be satisfied.

One embodiment satisfies 0.40 ≦ α ≦ 0.70.

In one embodiment, in the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is the I _ThMn12 , bcc- (Fe, Co, Ti) phase. When the maximum intensity of the peak caused by the 011 reflection is I _{bcc- (Fe, Co, Ti)} , I _{bcc- (Fe, Co, Ti)} / I _ThMn12 ≤ 0.75 is satisfied.

Certain embodiments, the powder X-ray diffraction pattern of the sintered body, a phase having the maximum intensity _{I ThMn12,} Th 2 _{Ni 17} type crystal structure peaks due to 002 reflecting the phase having the ThMn ₁₂ type crystal structure When the maximum intensity of the peak caused by the 023 reflection of No. 023 is I _Th2Ni17 , I _Th2Ni17 / I _ThMn12 ≦ 0.7 is satisfied.

Method for producing a rare earth magnet sintered body of the present disclosure, in an exemplary embodiment, the entire composition is represented by the following composition formula _{(1), R1 1-x} R2 x (Fe 1-y Co y) _{w-z Ti} z Cu _α O (1), R1 is Y or Y and Gd, Y is 50 mol% or more of the whole R1, and R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm. Sm is 50 mol% or more of the whole R2, and x, y, z, w, α, and β are 0.3 ≦ x ≦ 0.9, 0 ≦ y ≦ 0.4, 0, respectively. .38 ≦ z ≦ 0.70, 7 ≦ w ≦ 12, 0 ≦ α ≦ 0.70, 0.02 ≦ β ≦ 0.5, and 0 ≦ 1-x-2z / 3-0.092α-8β / A method for producing a sintered body for rare earth magnets, which comprises a phase having a ThMn _12- type crystal structure as a main phase, satisfying 15 ≦ 0.05, wherein the molten metal of the raw material is cooled to obtain an alloy, and the alloy. A step of crushing the fine powder to obtain a fine powder, a step of molding the fine powder to obtain a molded product, and a heat treatment of the molded product at 900 ° C. or higher and 1250 ° C. or lower and a pressure of 1000 MPa or lower for 5 minutes or longer and 50 hours or shorter to bake. Includes the step of obtaining a coalescence.

According to the embodiment of the present invention, it is possible to provide a sintered body for a rare earth magnet having few different phases that adversely affect the magnetic characteristics and a method for producing the same.

Sample No. It is a figure which shows the powder X-ray diffraction measurement result in 1-5. Sample No. It is a figure which shows the relative intensity of the bcc- (Fe, Co, Ti) phase obtained from the powder X-ray diffraction measurement result with respect to the value of 1-x-2z / 3-0.092α-8β / 15 in 1-5. is there. Sample No. It is a figure which shows the relative intensity of the 2-17 phase obtained from the powder X-ray diffraction measurement result with respect to the value of 1-x-2z / 3-0.092α-8β / 15 in 1-5. Sample No. It is a figure which shows the reflected electron image of the sample cross section in 1-5. Sample No. It is a figure which shows the powder X-ray diffraction measurement result in 6, 7 and 4. Sample No. It is a figure which shows the reflected electron image of the sample cross section in 7.

[Composition of sintered body for rare earth magnets]
The overall composition of the sintered body for rare earth magnets of the present disclosure is represented by the following composition formula (1).
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α O β (1)

Here, R1 is Y or Y and Gd, Y is 50 mol% or more of the whole R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm is Always included, Sm is 50 mol% or more of the whole R2. R1 is preferably Y only (excluding unavoidable impurities), and R2 is preferably Sm only (excluding unavoidable impurities). Further, x, y, z, w, α and β are 0.3 ≦ x ≦ 0.9, 0 ≦ y ≦ 0.4, 0.38 ≦ z ≦ 0.70, 7 ≦ w ≦ 12, respectively. Satisfy 0 ≦ α ≦ 0.70 and 0.02 ≦ β ≦ 0.5, and further satisfy the relational expression 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.05. Further, it is more preferable to satisfy the relational expression 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.03. Further, it is more preferable to satisfy 0.40 ≦ α ≦ 0.70.

As a result of diligent research by the present inventors, the bcc- (Fe, Co, Ti) phase, which adversely affects the magnetic properties, can be obtained by setting the sintered body in the composition range as shown in the above formula (1). , Th ₂ Ni ₁₇ type crystal or a compound having a similar structure thereof (hereinafter, may be referred to as 2-17 phase) can be suppressed.

[Reason for limitation of composition, etc.]
(Types of R1 and R2)
R1 is Y or Y and Gd, and Y is 50 mol% or more of the whole R1. When R1 is another element, a stable phase other than the 1-12 phase may be formed. For example, when R1 is Zr, a Th ₆ Mn ₂₃ type (hereinafter referred to as 6-23 phase) phase is generated, and a large amount of bcc- (Fe, Co, Ti) phase is also generated, so that a desired sintered body is produced. Cannot be obtained. Since Gd is expensive, it is preferable that R1 is only Y. Further, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr from the magnetic property values and phase stability of the 1-12 phase, and always contains Sm, and Sm is the entire R2. It is 50 mol% or more. From the viewpoint of the magnetic property value, it is more preferable that R2 is only Sm.

(R2 content)
The range of x (R2 substitution amount x) indicating the ratio of the atomic number of the content of R2 to the total amount of R1 and R2 is 0.3 ≦ x ≦ 0.9. If x is less than 0.3, the magnetic anisotropy of the 1-12 phase decreases, which is not preferable. On the other hand, if x is larger than 0.9, the stability of the 1-12 phase is lowered, and a bcc- (Fe, Co, Ti) phase or a 2-17 phase may be formed, which is not preferable.

(Ratio of Fe and Co)
The range of y (Co substitution amount y) indicating the ratio of the number of atoms of Co to the total of Fe and Co is 0 ≦ y ≦ 0.4. It is more preferable that y is 0.05 or more in order to avoid a possibility that the Curie temperature of the 1-12 phase is lowered. Further, if y is larger than 0.4, the volume magnetization and magnetic anisotropy magnetic field of the 1-12 phase decrease, which is not preferable.

(Ti content)
The range of z (Ti content z) indicating the atomic number ratio of the Ti content to the total amount of R1 and R2 is 0.38 ≦ z ≦ 0.70. If z is less than 0.38, the 2-17 phase and the bcc- (Fe, Co, Ti) phase are stably formed during sintering, which is not preferable. Further, if z is larger than 0.70, the magnetic properties of the 1-12 phase are lowered, which is not preferable. In order to obtain higher magnetic properties, particularly J _s , it is preferable that the amount of Ti is small. Specifically, it is more preferable that the range of z is 0.38 ≦ z ≦ 0.60. In addition, 50 mol% or less of Ti is an element that stabilizes the structure of 1-12 phase such as tungsten (W), vanadium (V), niobium (Nb), tantalum (Ta), molybdenum (Mo), and silicon (Si). It may be replaced.

(Cu content)
The range of α indicating the ratio of the number of atoms of the Cu content to the total amount of R1 and R2 is 0 ≦ α ≦ 0.7. When α is larger than 0.7, the ratio of the R—Cu phase as the sub-phase becomes high, the ratio of the main phase decreases, and the magnetization of the sintered body as a whole decreases, which is not preferable. Further, the range of α is more preferably 0.4 ≦ α ≦ 0.7. When α is smaller than 0.4, the amount of the liquid phase during the heat treatment is small, so that it becomes difficult to reduce the different phases during the solution treatment and to proceed with the densification during sintering.

(Total amount of Fe, Co, Ti)
The range of w indicating the atomic number ratio of the total amount of Fe, Co, and Ti to the total amount of R1 and R2 is 7 ≦ w ≦ 12. When w is larger than 12, the bcc- (Fe, Co, Ti) phase is remarkably generated, which is not preferable. Further, when w is smaller than 7, it is not preferable because a phase having a higher rare earth content than the 1-12 phase such as the 2-17 phase and adversely affecting the magnetic characteristics is remarkably generated.

(Oxygen content)
The range of 0.02 ≦ β ≦ 0.5 is appropriate for β, which indicates the atomic number ratio of the oxygen content to the total amount of R1 and R2. If β is less than 0.02, the fine powder before sintering tends to ignite, which makes handling difficult, which is not preferable. Further, when β is larger than 0.5, the ratio of the oxide phase in the sintered body becomes high, the ratio of the 1-12 phase decreases, and the magnetization of the magnet as a whole decreases, which is not preferable.

(Relationship between the amount of oxygen and the amount of other elements)
x, z, α, β satisfy the relational expression 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.05. Since the sintered body generally uses fine powder, the amount of oxygen is usually higher than that of the raw material alloy. Therefore, even if the alloy has few different phases at the stage of the raw material alloy, oxygen reacts with rare earths in the main phase and grain boundary phase (liquid phase at the time of sintering) at the time of fine grinding or sintering, and becomes an oxide phase. As a result, the 1-12 phase may be decomposed to form the bcc- (Fe, Co, Ti) phase. As a result of diligent research, the authors have found that Y contained in R1 is particularly susceptible to oxidation and how R1 is distributed to each phase. In the above relational expression, the amount of R1 consumed as the 1-12 phase from the value of z is 2z / 3, the amount of R1 consumed as the R—Cu phase from the value of α is 0.092α, and the value of β is the R oxide. R1 consumed as a phase is described as 8β / 15, respectively, and the upper and lower limits of the actual amount 1-x of R1 minus the calculated amount of R1 are defined. The smaller 1-x-2z / 3-0.092α-8β / 15 on the negative side, the more R1 required for the formation of the 1-12 phase, the R-Cu phase and the R oxide phase is insufficient. Especially, if it is less than 0, a large amount of bcc- (Fe, Co, Ti) phase is generated, which is not preferable. On the contrary, the larger the value on the plus side, the more R1 becomes surplus. Especially, if it is larger than 0.05, a large amount of rare earth content phase such as 2-17 phase is generated, which is not preferable. .. Further, when x, z, α, and β satisfy the relational expression 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.03, it is more than that of the 1-12 phase such as the 2-17 phase. It is preferable because it can suppress the formation of a phase having a high rare earth content. Further, as a simple method for examining the amount of the bcc- (Fe, Co, Ti) phase and the 2-17 phase present in the sintered body, powder X-ray diffraction measurement can be mentioned. Among the peaks of each phase, the phase having a ThMn ₁₂ type crystal structure (1-12 phase) has 002 reflection, the bcc- (Fe, Co, Ti) phase has 011 reflection, and the phase having a Th ₂ Ni ₁₇ type crystal structure (Th ₂ Ni ₁₇ type crystal structure). In the 2-17 phase), the peak caused by the 023 reflection is less affected by the other phases, and the peak intensity is high. Therefore, the maximum intensity of the peak caused by the 002 reflection of the 1-12 phase is I _ThMn12 , and the maximum intensity of the peak caused by the 011 reflection of the bcc- (Fe, Co, Ti) phase is I _{bcc- (Fe, Co, Ti). )} , When the maximum intensity of the peak caused by the 023 reflection of the 2-17 phase is I _Th2Ni17 , the relative intensities of the peaks of the bcc- (Fe, Co, Ti) phase and the 2-17 phase are I _bcc-, respectively _{. It} can be described as _{(Fe, Co, Ti)} / I _ThMn12 and I _Th2Ni17 / I _ThMn12 . It is desirable that the bcc- (Fe, Co, Ti) and 2-17 phases in the sintered body are as small as possible, and the relative strength of the bcc- (Fe, Co, Ti) phase is I _{bcc- (Fe, Co, Ti).} It is desirable that / I _ThMn12 is 0.75 or less. Further, it is desirable that the relative strength I _Th2Ni17 / I _{ThMn12 of} the 2-17 phase is 0.7 or less.

[Method of producing sintered body for rare earth magnets]
<Step A> Step of producing an alloy As a method for producing an alloy as a raw material for a sintered body for a rare earth magnet, a known method such as a mold casting method, a centrifugal casting method, a strip casting method, or a liquid ultra-quenching method is adopted. it can. In these methods, a molten alloy is prepared, and then the molten metal is cooled and solidified. It is desirable to suppress the formation of coarse bcc- (Fe, Co, Ti) phase and 2-17 phase as much as possible during solidification of the molten alloy. Coarse bcc- by adopting a method such as a strip casting method or a liquid ultra-quenching method, which has a relatively high cooling rate, to supply molten metal onto a rotating roll and solidify it to form a thin band or flaky alloy. The formation of (Fe, Co, Ti) phase and 2-17 phase can be suppressed. If the cooling rate during solidification is low, the size of the precipitated heterogeneous grains becomes large. When the grain size of the different phase contained in the alloy becomes large, it becomes difficult for the different phase to disappear during heat treatment such as a sintering step. The alloy heat treatment may be performed for the purpose of reducing the heterogeneous phase generated in the solidification process. The R—Cu phase melting point is 850 to 900 ° C., depending on the composition of the alloy. Therefore, the heat treatment temperature is preferably 900 ° C. or higher and 1250 ° C. or lower, and more preferably 1000 ° C. or higher and 1150 ° C. or lower. The heat treatment time depends on the heat treatment temperature, but is preferably 5 minutes or more and 50 hours or less. If the time is too short, there will not be enough reaction to eliminate the heterogeneity. If the time is too long, the rare earth elements will evaporate and oxidize, and the operational efficiency will be poor. Further, the alloy may be heat treated in hydrogen to introduce cracks prior to the grinding step. When Cu is contained, the R—Cu phase in the alloy can absorb and release hydrogen. According to this alloy, for example, hydrogen absorption occurs at a temperature of 250 ° C. to 400 ° C., and hydrogen release occurs between 540 ° C. and 660 ° C. Therefore, after the alloy is heated to 400 ° C. or higher in hydrogen to absorb hydrogen, the atmosphere can be switched to a vacuum atmosphere to sufficiently release hydrogen. In that case, the temperature for switching to the vacuum atmosphere is 700 ° C. or lower. As described above, the subphase contained in the present alloy absorbs and releases hydrogen at a temperature of at least 700 ° C. or lower. If the alloy is exposed to a hydrogen atmosphere at a temperature exceeding 700 ° C., decomposition of the main phase may occur due to a hydrogenation-disproportionation reaction. By absorbing and releasing hydrogen, the rare earth rich phase (secondary phase) undergoes volume expansion and contraction, and cracks occur between the main phase crystal grains and the subphase. This increases the crushing efficiency in the crushing process.

<Step B> Crushing step The alloy obtained in step A is crushed to obtain a fine powder. As a crushing method, a known method such as a jet mill, a stamp mill, or a ball mill can be adopted. Grinding is carried out in an inert gas such as nitrogen, argon or helium in order to suppress the oxidation of the powder and reduce the risk of ignition and explosion. A small amount of air or oxygen may be mixed with the inert gas in order to improve the handleability of the fine powder after pulverization. Considering the handling and moldability of the powder, the particle size of the fine powder after pulverization is D50 obtained by the laser diffraction method by the air flow dispersion method (volume-based median diameter of particles when the cumulative frequency reaches 50%). Is preferably 1 μm or more and 20 μm or less. If D50 is less than 1 μm, the risk of ignition increases and the mold is damaged during molding, which is not preferable. Further, if D50 is larger than 20 μm, densification is less likely to proceed in the sintering step, which is not preferable. The amount of oxygen in the sintered body is greatly affected by this pulverization step, and the pulverized particle size and the oxygen concentration in the pulverized gas greatly contribute. The finer the particle size of the fine powder and the higher the oxygen concentration in the pulverized gas, the larger the oxygen amount β in the sintered body. On the contrary, the coarser the particle size of the fine powder and the lower the oxygen concentration in the pulverized gas, the smaller the β value.

<Step C> Molding step The fine powder obtained in step B is molded to obtain a molded product. You may mold while applying a magnetic field at the time of molding to orient the crystal. Molding is a dry molding method in which dry fine powder is inserted into the cavity of the mold and molded, and a slurry (alloy powder is dispersed in the dispersion medium) is injected into the cavity of the mold to form the slurry. A known method including a wet molding method in which molding is performed while discharging the dispersion medium can be adopted.

<Step D> Sintering step A sintered body is obtained by heat-treating the molded product obtained in step C. As a sintering method, a method of holding at a high temperature in a vacuum or an inert gas atmosphere to proceed with solid phase sintering or liquid phase sintering, or a method of holding at a high temperature while applying pressure to a molded product. In order to prevent oxidation due to the atmosphere during sintering, it is preferable to perform sintering in a vacuum atmosphere or in an inert gas such as argon or helium. Further, since Sm evaporates remarkably at a high temperature, it is more preferable to suppress the evaporation of Sm by a method such as covering the molded product, sealing it, or sealing it together with a substance containing Sm. The sintering treatment temperature is 900 ° C. or higher and 1250 ° C. or lower. If the sintering treatment temperature is less than 900 ° C., a liquid phase is not sufficiently formed and it is difficult to densify. Further, if the sintering treatment temperature exceeds 1250 ° C., the 1-12 phase may be decomposed. The sintering treatment temperature is more preferably 1000 ° C. or higher and 1150 ° C. or lower. The sintering treatment time is 5 minutes or more and 50 hours or less. If the sintering process time is less than 5 minutes, densification may not proceed sufficiently. Further, if the sintering treatment time exceeds 50 hours, the lead time becomes long, which is not preferable in terms of operation. The pressure for pressure sintering is preferably 1000 MPa or less. Further, after the sintering step, heat treatment or diffusion treatment for the purpose of improving magnetic properties may be additionally performed.

<Experimental example>
Hereinafter, examples of the present invention will be specifically described, but the present invention is not limited to these examples.

First, a raw material alloy was produced by a strip casting method. The raw material metals of Y, Zr, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more are added to the evaporation of rare earth elements at the time of dissolution, and the final alloy composition obtained is the composition shown in Table 1. The target value was determined and weighed so as to be. The weighed metals were mixed and put into a silica crucible, and the temperature was raised to 1500 ° C. by high frequency induction heating to dissolve the raw materials. Then, the molten metal was cooled to 1450 ° C., temporarily stored in a tundish, and then supplied onto a copper cooling roll rotating at a peripheral speed of 1.5 m / s for cooling. The cooled alloy was crushed by a crusher installed under the cooling roll.

Each of the alloys having each composition obtained in the above step was weighed by 500 g and placed in a molybdenum container, and the container was inserted into a tubular heat treatment furnace. After replacing the inside of the furnace with Ar, heat treatment was performed at 1100 ° C. for 1.5 hours in an atmosphere in which Ar was perfuated at 2 L / min. After the heat treatment was completed, the heat treatment furnace was opened to cool the alloy. At this time, the average cooling rate from 1100 ° C. to 100 ° C. was 10 ° C./min or more.

The heat-treated alloy obtained in the above step was pulverized using a mortar in a glove box having an Ar flow atmosphere. The crushed powder was sieved with a 1 mm mesh, and the powder passed through the mesh was collected. Zinc stearate was added to the recovered pulverized powder, and the mixture was mixed with a locking mixer for 15 minutes. At this time, zinc stearate was added so that the weight ratio of the pulverized powder and zinc stearate was 100: 0.035.

The pulverized powder obtained in the above step was pulverized using an airflow jet mill PJM-100 manufactured by Nippon Pneumatic Industries Co., Ltd. to obtain fine powder. Nitrogen gas was used as the pulverizing gas, and pulverization was performed at a pulverization pressure of 7.5 MPa. The D50 of the fine powder at this time was 5 μm in each case.

The fine powder obtained in the above step was molded in a glove box having an Ar flow atmosphere. A hand press was used for molding, and a cylindrical molded body having a diameter of 16 mm and a height of 20 mm was produced. After molding, the iron capsule was filled with the molded product. The material of the iron capsule was S20C, the inner diameter was 16 mm, the height was 20 mm, and the thickness was 2 mm.

The iron capsule filled with the molded body was subjected to electron beam welding in vacuum, and the capsule container and lid were welded to seal the capsule.

The sealed sample was subjected to hot isotropic pressurization (HIP) treatment. Argon was used as the pressure medium gas, and the treatment was performed at a gas pressure of 180 MPa. The temperature was 1100 ° C. and the holding time was 3 hours.

The sample obtained in the above step was cut with an outer peripheral blade cutting machine, and the HIP body contained in the capsule was taken out. A part of the HIP body was crushed in a mortar and classified using a 425 μm mesh and a 75 μm mesh.

Using crushed powder with a particle size of 75 to 425 μm, component analysis of Y, Zr, Sm, Fe, Co, Ti, and Cu by ICP (inductively coupled plasma) emission spectroscopy and carbon by combustion and infrared absorption methods Quantitative analysis was performed. The amount of oxygen and the amount of nitrogen were analyzed by the Inert gas melting / heat conduction method using pulverized powder having a particle size of 425 μm or more. Further, the carbon content was analyzed by a combustion / infrared absorption method using pulverized powder having a particle size of 75 to 425 μm. From the analysis results, the values of x, y, z, w, α, β, and 1-x-2z / 3-0.092α-8β / 15 of each sintered body were obtained.

Powder X-ray diffraction was performed using pulverized powder having a particle size of less than 75 μm. A Bragg-Brentano focused beam type wide-angle X-ray diffractometer (X-ray diffuser, XRD, D8 ADVANCED / TXS manufactured by Bruker AX Co., Ltd.) was used as the apparatus. A Cu rotating anti-cathode was used as the X-ray generation source, the applied voltage was 45 kV, the current was 360 mA, and Ni was used as the Kβ filter. The scanning axis was 2θ / θ interlocking operation, the interval was 0.04 °, the speed was 0.6s / step, and the range of 20 ° ≤ 2θ ≤ 80 ° was scanned at room temperature. From the X-ray intensity profile, the maximum intensity of the peak due to the 002 reflection of the 1-12 phase is I _ThMn12 , and the maximum intensity of the peak due to the 011 reflection of the bcc- (Fe, Co, Ti) phase is I _{α- (Fe). , Co, Ti)} , the maximum intensity of the peak caused by the 023 reflection of the 2-17 phase is I _Th2Ni17, and the relative X-ray intensity I _{bcc- (Fe, Co,} Ti) of the bcc- (Fe, Co, Ti) phase _{. Ti)} / I _ThMn12 and the relative X-ray intensities I _Th2Ni17 / I _ThMn12 of the 2-17 phase were determined, respectively.

The cut HIP body was embedded in a resin, polished, the cross section of the HIP body was observed with a scanning electron microscope (SEM), and a local composition analysis was performed by EDX. As the SEM, JCM-6000 Plus NeoScope (registered trademark) manufactured by JEOL Ltd. was used, and a reflected electron image was acquired and EDX analysis was performed at an accelerating voltage of 15 kV.

Table 1 shows the compositions of the prepared HIPs, with values of x, y, z, w, α, β, 1-x-2z / 3-0.092α-8β / 15, and I _{bcc- (Fe, Co,} The values of _Ti) / I _ThMn12 and I _Th2Ni17 / I _ThMn12 are shown in Table 2.

No. 1 to 5 are experimental examples in which the Y and Sm contents are changed. In all the experimental examples, the values of x, y, z, w, α, and β are all within the preferable range. No. The powder XRD pattern of the samples 1 to 5 is shown in FIG. In addition, No. _{Values of} I _{bcc- (Fe, Co, Ti)} / I _ThMn12 and I _Th2Ni17 / I _ThMn12 and 1-x-2z / 3-0.092α-8β / 15 obtained from the powder XRD patterns of the samples 1 to 5 The relationship between the values of is shown in FIGS. 2 and 3. No. 1 in which the value of 1-x-2z / 3-0.092α-8β / 15 is less than 0. A large amount of bcc- (Fe, Co, Ti) phase was present in the samples 1 and 2, and the value of I _{bcc- (Fe, Co, Ti)} / I _ThMn12 was high. No. 1 in which the value of 1-x-2z / 3-0.092α-8β / 15 is in the range of 0 or more and 0.05 or less. The values of I _{bcc- (Fe, Co, Ti)} / I _ThMn12 and I _Th2Ni17 / I _ThMn12 were suppressed in the samples 3 and 4 because the formation of the bcc- (Fe, Co, Ti) phase and the 2-17 phase was suppressed. Was a low result. In particular, No. 1 in which the value of 1-x-2z / 3-0.092α-8β / 15 is in the range of 0 or more and 0.03 or less. Sample 3 had very low values of I _{bcc- (Fe, Co, Ti)} / I _ThMn12 and I _Th2Ni17 / I _ThMn12 . No. 1-x-2z / 3-0.092α-8β / 15 with a value greater than 0.05. In the sample _No. 5, a large amount of 2-17 phases were present, and the value of I _Th2Ni17 / I _ThMn12 was high. No. The reflected electron images of the sample cross sections 1 to 5 are shown in FIG. No. A large amount of bcc- (Fe, Co, Ti) phase was observed in sample 1. In addition, No. A large amount of 2-17 phases were observed in sample 5. No. In the samples 3 and 4, the formation of the bcc- (Fe, Co, Ti) phase and the 2-17 phase was suppressed, which was in good agreement with the result of powder XRD.

No. 6, 7 and No. The powder XRD pattern of the sample of No. 4 is shown in FIG. No. In the sample No. 6, Cu was not added, and No. It is a sample aiming at a composition having a value of 1-x-2z / 3-0.092α-8β / 15 equivalent to that of sample 4. If the value of 1-x-2z / 3-0.092α-8β / 15 is within the preferable range regardless of the presence or absence of Cu, the bcc- (Fe, Co, Ti) phase and the 2-17 phase are generated. It was possible to suppress, and the values of I _{bcc- (Fe, Co, Ti)} / I _ThMn12 and I _Th2Ni17 / I _ThMn12 were low. In addition, No. The sample of No. 7 is No. This is a sample aimed at a composition in which Y of sample 4 is replaced with Zr. No. A large amount of bcc- (Fe, Co, Ti) phase was present in the sample _{No. 7 ,} and I _{bcc- (Fe, Co, Ti)} / I _ThMn12 was very high. No. The reflected electron image of the sample cross section of No. 7 is shown in FIG. No. In the sample No. 7, there are 6-23 phases, and from the results of EDX point analysis, Zr is contained in about 13% by 13 atoms and Ti is contained in about 12% in this 6-23 phase, which is compared with the overall composition. Zr and Ti were very concentrated.

The sintered body for rare earth magnets of the present disclosure can be used for rare earth magnets.

Claims (6)

The overall composition is represented by the following composition formula (1).
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α O β (1)
R1 is Y or Y and Gd, and Y is 50 mol% or more of the total R1.
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is 50 mol% or more of the total R2.
x, y, z, w, α, and β are, respectively.
0.3 ≤ x ≤ 0.9,
0 ≦ y ≦ 0.4,
0.38 ≤ z ≤ 0.70,
7 ≦ w ≦ 12,
0 ≤ α ≤ 0.70,
0.02 ≤ β ≤ 0.5, and 0 ≤ 1-x-2z / 3-0.092α-8β / 15 ≤ 0.05
A sintered body for rare earth magnets having a phase having a ThMn _12- type crystal structure as a main phase. 0 ≦ 1-x-2z / 3-0.092α-8β / 15 ≦ 0.03
The sintered body for rare earth magnets according to claim 1, which satisfies the above requirements. 0.40 ≤ α ≤ 0.70
The sintered body for rare earth magnets according to claim 1 or 2, which satisfies the above requirements. In the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is caused by the 011 reflection of the _IThMn12 , bcc- (Fe, Co, Ti) phase. When the maximum intensity of the peak is I _{bcc- (Fe, Co, Ti)} ,
I _{bcc- (Fe, Co, Ti)} / I _ThMn12 ≤ 0.75
The sintered body for rare earth magnets according to any one of claims 1 to 3, which satisfies the above. In the powder X-ray diffraction pattern of the sintered body, due to the maximum intensity of the peak attributable to the 002 reflection of the phases with the ThMn ₁₂ type crystal structure 023 reflecting the phase with _{I ThMn12,} Th 2 _{Ni 17} type crystal structure When the maximum intensity of the peak to be _crystallized is I _Th2Ni17 ,
I _Th2Ni17 / I _ThMn12 ≤ 0.7
The sintered body for rare earth magnets according to any one of claims 1 to 4, which satisfies the above requirements. The overall composition is represented by the following composition formula (1).
_{R1 1-x R2 x (Fe} 1-y Co y) w-z Ti z Cu α O β (1)
R1 is Y or Y and Gd, and Y is 50 mol% or more of the total R1.
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is 50 mol% or more of the total R2.
x, y, z, w, α, and β are, respectively.
0.3 ≤ x ≤ 0.9,
0 ≦ y ≦ 0.4,
0.38 ≤ z ≤ 0.70,
7 ≦ w ≦ 12,
0 ≤ α ≤ 0.70,
0.02 ≤ β ≤ 0.5, and 0 ≤ 1-x-2z / 3-0.092α-8β / 15 ≤ 0.05
A method for producing a sintered body for a rare earth magnet, which comprises a phase having a ThMn _12- type crystal structure as a main phase, which satisfies the above.
The process of cooling the molten metal of the raw material to obtain an alloy,
The process of crushing the alloy to obtain fine powder and
The process of molding the fine powder to obtain a molded product,
A method for producing a sintered body for a rare earth magnet, which comprises a step of heat-treating the molded body at 900 ° C. or higher and 1250 ° C. or lower and a pressure of 1000 MPa or lower for 5 minutes or more and 50 hours or less to obtain a sintered body.