JP7196666B2 - Sintered body for rare earth magnet and method for producing the same - Google Patents

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 machinery, and home appliances.

Nd--Fe--B system magnets can be cited as typical high-performance magnets. Nd--Fe--B magnets are mainly used in drive motors for electric vehicles (EV, HV, PHV, etc.) and hybrid vehicles. As the need for motors with higher efficiency and smaller size increases, the development of permanent magnets with higher magnetic properties is expected.

An RT ₁₂ compound having a ThMn ₁₂ type crystal structure or a structure similar thereto is attracting attention as one of the candidates for main phase system alloys for permanent magnets that surpasses the magnetic properties of Nd--Fe--B magnets. The RT ₁₂ compound is R ₂ T ₁₄ B (R is at least one rare earth element, T is at least one iron group transition metal element containing at least Fe), which is a compound that constitutes the main phase of an Nd—Fe—B magnet. ) Higher magnetic properties are expected due to higher concentrations of iron group transition metals. Hereinafter, a phase composed of an RT ₁₂ compound having a ThMn ₁₂ -type crystal structure or a similar structure may be referred to as a 1-12 phase.

Patent Document 1 discloses a rare earth permanent magnet in which a portion of Fe, which is a T element, is partially replaced with Ti, which is a structure stabilizing element, to improve thermal stability in exchange for high magnetization. there is

In Patent Document 2, by partially substituting the R element of the RFe ₁₂ -based compound with a substituting element such as Zr or Hf, the amount of the substituting element such as Ti substituting the transition metal element is reduced to increase the saturation magnetization. Rare earth permanent magnets are disclosed which are stabilized while preserving the ThMn type ₁₂ crystal structure.

Further, Patent Document 3 discloses an R'--Fe--Co system ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R'--Fe--Co system ferromagnetic It is described that the alloy exhibits high magnetic properties due to having a ThMn ₁₂ -type crystal structure produced by an ultraquenching method.

Further, Patent Document 4 describes that by adding Cu, a non-magnetic and low melting point 1-4 composition (SmCu ₄ phase) phase is generated, and sintering and high coercive force are possible. .

Further, in Patent Document 5, at least one of Sm ₅ Fe ₁₇ -based phase, SmCo ₅ -based phase, Sm ₂ O ₃ -based phase, and Sm ₇ Cu ₃ -based phase is added as a secondary phase to a ThMn ₁₂ -type main phase. It is described that the coercive force can be increased by including the element.

Further, Patent Document 6 describes that the addition of Cu generates a liquid phase and enables the formation of a dense bulk body.

In addition, in Patent Document 7, by producing a ferromagnetic alloy whose main phase is a ThMn ₁₂ type phase containing Y by a strip casting method, the main phase composition is less uniform and the main phase ratio is high. is obtained.

Further, Patent Document 8 describes that a magnetic material having a Y-containing ThMn ₁₂ -type phase as a main phase can provide high saturation magnetization and anisotropic magnetic field.

JP-A-64-76703 JP-A-4-322406 JP 2015-156436 A Japanese Patent Application Laid-Open No. 2001-189206 Japanese Unexamined Patent Application Publication No. 2017-112300 WO2016/162990 JP 2018-103211 A JP 2018-125512 A

As one of the conditions for a sintered body to be used for a high-performance magnet, it is necessary to have a structure with few heterogeneous phases that adversely affect magnetic properties. When a soft magnetic phase typified by the bcc-Fe phase exists in the sintered body, the soft magnetic phase becomes a starting point for magnetization reversal, and magnetization reversal easily progresses. Magnetic properties are significantly degraded. Therefore, a sintered body in which such a heterogeneous phase of soft magnetism does not exist as much as possible is required.

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

On the other hand, in the rare earth permanent magnet described in Patent Document 2, the _ThMn12 structure is stabilized by substituting a transition metal element with Ti or the like, but the effect is not necessarily sufficient.

The R'--Fe--Co ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structure stabilizing element M (such as Ti), so it exhibits high magnetization, large magnetic anisotropy, and a high Curie temperature. However, due to the non-equilibrium phase, the main phase compound may decompose in a high temperature densification process such as sintering.

The rare earth magnet described in Patent Document 4 may not have high magnetic physical properties due to the large amount of Ti added.

In the rare earth magnet described in Patent Document 5, when the rare earth-rich secondary phase Sm ₇ Cu ₃ is used, a phase richer in rare earth than the main phase may be generated due to the reaction between the main phase and Sm ₇ Cu ₃ during heat treatment. 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 a high Curie temperature can be obtained, and the density as a bulk body is high. Since it is a non-equilibrium phase, the main phase compound may decompose in a high temperature process such as sintering at 1000° C. or higher.

Since the composition of the ferromagnetic alloy described in Patent Document 7 and the magnet material described in Patent Document 8 does not take into account the influence of oxygen that is inevitably mixed in the production process of the sintered body, oxygen is preferred as a rare earth element. It is feared that a soft magnetic phase such as a bcc-Fe phase may be generated due to the decomposition of the main phase.

Embodiments of the present disclosure provide a sintered body for a rare earth magnet with less heterogeneous phases that adversely affect magnetic properties, and a method for producing the same.

In an exemplary embodiment, the overall composition of the sintered body for a rare earth magnet whose main phase is a phase having a ThMn ₁₂ -type crystal structure of the present disclosure is represented by the following composition formula (1), and R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α O _β (1), R1 is Y or Y and Gd, Y is 50 mol % or more of the total of R1, R2 is Sm, La, is at least one selected from the group consisting of Ce, Nd and Pr, must contain Sm, Sm is 50 mol% or more of the total R2, and x, y, z, w, α, and β are each 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.

Some embodiments have 0≦1−x−2z/3−0.092α−8β/15≦0.03
satisfy.

Some embodiments satisfy 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 I _ThMn12 , bcc-(Fe, Co, Ti) phase where _{Ibcc- ( Fe,Co,Ti)} is the maximum intensity of the peak due to the 011 reflection of .

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 I _ThMn12 , the phase having the Th ₂ Ni ₁₇ type crystal structure where I _Th2Ni17 is the maximum intensity of the peak due to the 023 reflection of , I _Th2Ni17 /I _ThMn12 ≦0.7 is satisfied.

In an exemplary embodiment of the method for producing a sintered body for a rare earth magnet of the present disclosure, the overall composition is represented by the following composition formula (1), R1 _1-x R2 _x (Fe _1-y Co _y ) _{wzTizCuαO _ _} (1), R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, 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/ A method for producing a sintered body for a rare earth magnet, the main phase being a phase having a ThMn type ₁₂ crystal structure, which satisfies 15≦0.05, comprising the steps of: obtaining an alloy by cooling a molten raw material; A step of pulverizing the fine powder to obtain a fine powder, a step of molding the fine powder to obtain a molded body, and a heat treatment of 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. and obtaining a body.

According to the embodiments of the present invention, it is possible to provide a sintered body for a rare earth magnet with less heterogeneous phases that adversely affect magnetic properties, and a method for producing the same.

Sample no. 1 is a diagram showing powder X-ray diffraction measurement results in 1 to 5. FIG. Sample no. 1 to 5, the relative intensity of the bcc-(Fe, Co, Ti) phase obtained from the powder X-ray diffraction measurement results for the values of 1-x-2z/3-0.092α-8β/15. be. Sample no. 1 is a diagram showing the relative intensity of the 2-17 phase obtained from powder X-ray diffraction measurement results with respect to the values of 1-x-2z/3-0.092α-8β/15 in 1 to 5. FIG. Sample no. 1 is a diagram showing backscattered electron images of cross sections of samples in 1 to 5. FIG. Sample no. It is a figure which shows the powder X-ray-diffraction measurement result in 6, 7 and 4. Sample no. 7 is a diagram showing a backscattered electron image of a sample cross section in 7. FIG.

[Composition of sintered body for rare earth magnet]
The overall composition of the sintered body for a rare earth magnet of the present disclosure is represented by the following compositional 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 total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm Sm must be included, and Sm is 50 mol % or more of the entire R2. R1 is preferably Y only (excluding unavoidable impurities), and R2 is preferably Sm only (excluding unavoidable impurities). 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, It satisfies 0≤α≤0.70 and 0.02≤β≤0.5, and further satisfies the relational expression 0≤1-x-2z/3-0.092α-8β/15≤0.05. More preferably, the relational expression 0≤1-x-2z/3-0.092α-8β/15≤0.03 is satisfied. Moreover, it is more preferable to satisfy 0.40≦α≦0.70.

As a result of intensive research by the present inventors, by setting the composition range of the sintered body as shown in the above formula (1), the bcc-(Fe, Co, Ti) phase, which adversely affects the magnetic properties, and , Th ₂ Ni ₁₇ -type crystals or a compound phase having a similar structure (hereinafter sometimes referred to as a 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 total of R1. When R1 is another element, a stable phase other than the 1-12 phase may be produced. For example, when R1 is Zr, a Th ₆ Mn ₂₃ type (hereinafter referred to as a 6-23 phase) phase is generated, and a large amount of the bcc-(Fe, Co, Ti) phase is also generated. is not obtained. Since Gd is expensive, it is preferable that R1 is only Y. In addition, from the magnetic physical property value and phase stability of the 1-12 phase, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is the total of R2 50 mol % or more. From the viewpoint of magnetic physical properties, R2 is more preferably Sm only.

(Content of R2)
The range of x (R2 substitution amount x), which indicates the atomic number ratio 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 is lowered, which is not preferable. On the other hand, if x is more than 0.9, the stability of the 1-12 phase is lowered, and the bcc-(Fe, Co, Ti) phase and the 2-17 phase may be generated, which is not preferable.

(ratio of Fe and Co)
The range of y (Co substitution amount y), which indicates the atomic number ratio of Co to the sum of Fe and Co, is 0≦y≦0.4. In order to avoid the risk of lowering the Curie temperature of the 1-12 phase, y is more preferably 0.05 or more. Also, if y is larger than 0.4, the volume magnetization and magnetic anisotropic magnetic field of the 1-12 phase are lowered, which is not preferable.

(Content of Ti)
The range of z (Ti content z), which indicates the atomic number ratio of the Ti content to the total amount of R1 and R2, is 0.38≦z≦0.70. When z is less than 0.38, the 2-17 phase and the bcc-(Fe, Co, Ti) phase are stably generated during sintering, which is not preferable. Moreover, 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, especially _Js , the smaller the Ti content, the better. Specifically, it is more preferable that the range of z is 0.38≦z≦0.60. Note that 50 mol% or less of Ti is an element that stabilizes the 1-12 phase structure, such as tungsten (W), vanadium (V), niobium (Nb), tantalum (Ta), molybdenum (Mo), and silicon (Si). may be replaced.

(Cu content)
The range of α, which indicates the atomic number ratio of the Cu content to the total amount of R1 and R2, is 0≦α≦0.7. If α is larger than 0.7, the ratio of the R—Cu phase, which is the subphase, increases, the ratio of the main phase decreases, and the magnetization of the sintered body as a whole decreases, which is not preferable. More preferably, the range of α is 0.4≦α≦0.7. If α is less than 0.4, the amount of liquid phase during heat treatment is reduced, so that the reduction of heterogeneous phases during solution treatment and the densification during sintering are difficult to progress.

(Total amount of Fe, Co and Ti)
The range of w, which indicates 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 produced significantly, which is not preferable. If w is less than 7, a phase such as the 2-17 phase, which has a higher rare earth element content than the 1-12 phase and has an adverse effect on the magnetic properties, is not preferable.

(Oxygen content)
β, which indicates the atomic number ratio of the oxygen content to the total amount of R1 and R2, is suitably in the range of 0.02≦β≦0.5. If β is less than 0.02, the fine powder before sintering tends to ignite, making handling difficult. If β is more than 0.5, the ratio of the oxide phase in the sintered body increases, 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, and β satisfy the relational expression 0≤1-x-2z/3-0.092α-8β/15≤0.05. Since the sintered body generally uses fine powder, it usually has a higher oxygen content than the raw material alloy. Therefore, even in alloys with few heterogeneous phases at the raw material alloy stage, oxygen reacts with rare earth elements in the main phase and grain boundary phase (liquid phase during sintering) during pulverization and sintering to form oxide phases. As the 1-12 phase decomposes, the bcc-(Fe, Co, Ti) phase may be generated. As a result of intensive research, the authors found out that Y contained in R1 is particularly easily oxidized and how R1 is distributed in each phase. From the value of z, the amount of R1 consumed as the 1-12 phase is 2z/3, from the value of α, the amount of R1 consumed as the R-Cu phase is 0.092α, and from the value of β, the R oxide It is a formula that describes the R1 consumed as a phase as 8β/15, respectively, and defines the upper and lower limits of the actual amount of R1 1−x minus the amount of R1 calculated from z, α, and β. As 1-x-2z/3-0.092α-8β/15 becomes smaller on the minus side, it means that R necessary for generating 1-12 phase, R-Cu phase and R oxide phase is insufficient. In particular, when it is less than 0, a large amount of bcc-(Fe, Co, Ti) phase is generated, which is not preferable. Conversely, the larger the value on the plus side, the more R1 becomes excessive, and especially when it is larger than 0.05, it is not preferable because a large amount of phases such as 2-17 phases having a higher rare earth content than 1-12 phases are generated. . Also, when x, z, α, and β satisfy the relational expression 0 ≤ 1-x-2z/3-0.092α-8β/15 ≤ 0.03, more than 1-12 phases such as 2-17 phases This is preferable because it can suppress the formation of a phase with a high rare earth content. Moreover, powder X-ray diffraction measurement can be mentioned as a simple method for examining the extent to which the bcc-(Fe, Co, Ti) phase and 2-17 phase are present in the sintered body. 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 Th ₂ Ni ₁₇ type crystal structure ( 2-17 phase), the peak due to 023 reflection is less affected by other phases, and the peak intensity is high. Therefore, 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 _{bcc-(Fe, Co, Ti )} , the relative intensities of the peaks of the _bcc- (Fe, Co, Ti) phase and the 2-17 phase are I _{bcc- (Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 . It is desirable that the bcc-(Fe, Co, Ti) phase and the 2-17 phase in the sintered body are as small as possible, and the relative strength of the bcc-(Fe, Co, Ti) phase I _{bcc-(Fe, Co, Ti)} /I _ThMn12 is preferably 0.75 or less. Also, the relative intensity I _Th2Ni17 /I _ThMn12 of the 2-17 phase is preferably 0.7 or less.

[Method for producing sintered body for rare earth magnet]
<Step A> A step of producing an alloy As a method for producing an alloy that is a raw material for a sintered body for a rare earth magnet, a known method such as a die casting method, a centrifugal casting method, a strip casting method, or a liquid ultra-quenching method is adopted. can. These methods involve making a molten alloy and then cooling and solidifying the molten alloy. It is desirable to minimize the formation of coarse bcc-(Fe, Co, Ti) phases and 2-17 phases during solidification of the molten alloy. Coarse bcc- It is possible to suppress the generation of (Fe, Co, Ti) phases and 2-17 phases. When the cooling rate during solidification is low, the grain size of the precipitated heterogeneous phase becomes large. When the grain size of the hetero-phase contained in the alloy becomes large, it becomes difficult to eliminate the hetero-phase during heat treatment such as a sintering process. An alloy heat treatment may be performed for the purpose of reducing heterogeneous phases generated during the solidification process. Depending on the composition of the alloy, the R—Cu phase melting point is 850-900°C. Therefore, the heat treatment temperature is preferably 900° C. or higher and 1250° C. or lower, more preferably 1000° C. or higher and 1150° C. or lower. Moreover, although the heat treatment time depends on the heat treatment temperature, it is desirable that the heat treatment time is 5 minutes or more and 50 hours or less. If the time is too short, there will not be enough reaction to eliminate the heterophases. If the time is too long, evaporation and oxidation of the rare earth elements will occur, and the operational efficiency will be poor. Additionally, prior to the grinding step, the alloy may be heat treated in hydrogen to introduce cracks. When containing Cu, the R—Cu phase in the alloy can absorb and release hydrogen. With this alloy, for example, hydrogen absorption occurs at temperatures between 250°C and 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. Thus, the subphase contained in the present alloy undergoes hydrogen absorption and desorption at a temperature of at least 700° C. or less. If the alloy is exposed to a hydrogen atmosphere at a temperature above 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 (subphase) expands and contracts in volume, and cracks occur between the main phase grains and the subphase. This increases the pulverization efficiency in the pulverization process.

<Step B> Grinding Step The alloy obtained in Step A is pulverized to obtain a fine powder. As a pulverization method, a known method such as a jet mill, a stamp mill, or a ball mill can be used. Grinding is performed in an inert gas such as nitrogen, argon, or helium to suppress oxidation of the powder and reduce the risk of ignition or 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 (the volume-based median diameter of the particles when the cumulative frequency is 50%) obtained by the laser diffraction method by the air dispersion method. 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, when D50 is larger than 20 μm, densification is difficult to proceed in the sintering process, which is not preferable. The amount of oxygen in the sintered body is greatly affected by the pulverization process, and the grain size of the pulverization and the oxygen concentration in the pulverization gas greatly contribute. The oxygen content β in the sintered body increases as the particle size of the fine powder becomes finer and as the oxygen concentration in the pulverizing gas increases. Conversely, the larger the particle size of the fine powder and the lower the oxygen concentration in the pulverization gas, the smaller the value of β.

<Step C> Molding Step The fine powder obtained in Step B is molded to obtain a compact. In order to orient the crystals, molding may be performed while applying a magnetic field during molding. In addition, the molding method is a dry molding method in which dry fine powder is inserted into the mold cavity and molded, and slurry (alloy powder is dispersed in the dispersion medium) is injected into the mold cavity and A known method including a wet molding method in which molding is performed while discharging a dispersion medium can be employed.

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

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

First, a raw material alloy was produced by a strip casting method. Raw material metals of Y, Zr, Sm, Fe, Co, Ti, and Cu with a purity of 99.9% or more are combined with the evaporation of rare earth elements during melting, and the final alloy composition obtained is the composition shown in Table 1. A target value was determined and weighed so that The weighed metals were mixed and charged into a silica crucible, and heated to 1500° C. by high-frequency induction heating to melt the raw materials. Thereafter, the molten metal was cooled to 1450° C., temporarily stored in a tundish, and then fed onto a cooling roll made of copper rotating at a peripheral speed of 1.5 m/s for cooling. The cooled alloy was pulverized by a pulverizer installed below the cooling rolls.

500 g of the alloy having each composition obtained in the above steps was weighed and placed in a molybdenum container, and the container was inserted into a tubular heat treatment furnace. After replacing the atmosphere in the furnace with Ar, heat treatment was performed at 1100° C. for 1.5 hours in an atmosphere in which 2 L/min of Ar was allowed to flow. 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 in an Ar atmosphere. The pulverized powder was sieved through a 1 mm mesh, and the powder that passed through the mesh was recovered. Zinc stearate was added to the collected pulverized powder and mixed with a rocking mixer for 15 minutes. At this time, zinc stearate was added so that the weight ratio of ground powder and zinc stearate was 100:0.035.

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

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

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

The sealed samples were subjected to hot isostatic pressing (HIP) treatment. Argon was used as a 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 a peripheral blade cutting machine, and the HIP body in the capsule was taken out. A portion of the HIP body was pulverized in a mortar and classified using 425 μm mesh and 75 μm mesh.

Using pulverized powder with a particle size of 75 to 425 μm, the components of Y, Zr, Sm, Fe, Co, Ti, and Cu are analyzed by ICP (inductively coupled plasma) emission spectrometry, and carbon is analyzed by combustion and infrared absorption. Quantitative analysis was performed. Using pulverized powder with a particle size of 425 μm or more, the oxygen content and nitrogen content were analyzed by the inert gas fusion/heat conduction method. Also, using pulverized powder with a particle size of 75 to 425 μm, the carbon content was analyzed by the combustion/infrared absorption method. 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 determined.

Powder X-ray diffraction was performed using pulverized powder having a particle size of less than 75 μm. As an apparatus, a Bragg-Brentano concentrated beam type wide-angle X-ray diffractometer (X-ray diffractiometer, XRD, D8 ADVANCED/TXS manufactured by Bruker AX Co., Ltd.) was used. A Cu rotating anticathode was used as the X-ray source, the applied voltage was 45 kV, the current was 360 mA, and the Kβ filter was Ni. The scanning axis was interlocked with 2θ/θ, the interval was 0.04°, the speed was 0.6 s/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 due to the 023 reflection of the 2-17 phase is I _Th2Ni17 , and the relative X-ray intensity of the bcc-(Fe, Co, Ti) phase is I _bcc- (Fe, Co, Ti). _Ti) 2 /I _ThMn12 and the relative X-ray intensity I _Th2Ni17 /I _ThMn12 of the 2-17 phase were determined respectively.

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

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

No. 1 to 5 are experimental examples in which the contents of Y and Sm were changed. All the values of x, y, z, w, α, and β are within preferable ranges in all experimental examples. No. The powder XRD patterns of samples 1-5 are shown in FIG. Also, No. Values of I _{bcc-(Fe,Co,Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 and 1-x-2z/3-0.092α-8β/15 determined from powder XRD patterns of samples 1-5 are shown in FIGS. 2 and 3. FIG. No. 1-x-2z/3-0.092α-8β/15 having a value of less than 0. A large amount of bcc-(Fe,Co,Ti) phase was present in samples 1 and 2, resulting in high values of Ibcc- _(Fe,Co,Ti) / _IThMn12 . No. 1-x-2z/3-0.092α-8β/15 in the range of 0 or more and 0.05 or less. For samples 3 and 4, the values of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 were suppressed due to suppression of bcc-(Fe, Co, Ti) and 2-17 phases. gave a low result. In particular, No. 1 having a value of 1-x-2z/3-0.092α-8β/15 in the range of 0 or more and 0.03 or less. Sample No. 3 resulted in very low values for both I bcc - _{(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 . No. 1-x-2z/3-0.092α-8β/15 value greater than 0.05. Sample No. 5 contained a large amount of 2-17 phase, resulting in a high value of I _Th2Ni17 /I _ThMn12 . No. Backscattered electron images of cross sections of samples 1 to 5 are shown in FIG. No. A large amount of bcc-(Fe, Co, Ti) phase was observed in sample 1. Also, No. A large amount of 2-17 phase was observed in sample 5. No. In samples 3 and 4, the formation of the bcc-(Fe, Co, Ti) phase and the 2-17 phase was suppressed, and the results corresponded well with the powder XRD results.

No. 6, 7 and no. The powder XRD pattern of sample No. 4 is shown in FIG. No. No. 6 sample does not add Cu. This is a sample aiming at a composition with a value of 1-x-2z/3-0.092α-8β/15, which is equivalent to the sample No. 4. With or without Cu, if the value of 1-x-2z/3-0.092α-8β/15 is within the preferred range, the formation of the bcc-(Fe, Co, Ti) phase and the 2-17 phase is can be suppressed, resulting in low values of I bcc - _{(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 . Also, No. 7 sample is No. This is a sample aiming at a composition in which Y of the sample of No. 4 is replaced with Zr. No. A large amount of bcc-(Fe,Co,Ti) phase was present in sample No. 7, resulting in very high Ibcc- _(Fe,Co,Ti) / _IThMn12 . No. FIG. 6 shows a backscattered electron image of the sample cross section of No. 7. No. The 6-23 phase exists in sample 7, and from the EDX point analysis results, this 6-23 phase contains about 13 atomic % of Zr and about 12 atomic % of Ti. 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 compositional formula (1),
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α O _β (1)
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of 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 of R2,
x, y, z, w, α, and β are each:
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 a rare earth magnet having a phase having a ThMn type ₁₂ crystal structure as a main phase, satisfying the above. 0≤1-x-2z/3-0.092α-8β/15≤0.03
The sintered body for a rare earth magnet according to claim 1, which satisfies 0.40≤α≤0.70
3. The sintered body for a rare earth magnet according to claim 1, which satisfies 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 I _ThMn12 , due to the 011 reflection of the bcc-(Fe, Co, Ti) phase. When the maximum intensity of the peak to be _{Ibcc-(Fe, Co, Ti)} is
Ibcc- _{(Fe,Co,Ti) /IThMn12≤0.75}
4. The sintered body for a rare earth magnet according to any one of claims 1 to 3, which satisfies 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 I _ThMn12 , and the maximum intensity of the peak due to the 023 reflection of the phase having the Th ₂ Ni ₁₇ type crystal structure When I _Th2Ni17 is the maximum intensity of the peak that
I _Th2Ni17 /I _ThMn12 ≦0.7
5. The sintered body for a rare earth magnet according to any one of claims 1 to 4, which satisfies The overall composition is represented by the following compositional formula (1),
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α O _β (1)
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of 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 of R2,
x, y, z, w, α, and β are each:
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 having a phase having a ThMn type ₁₂ crystal structure as a main phase, satisfying
a step of cooling the raw material molten metal to obtain an alloy;
pulverizing the alloy to obtain a fine powder;
a step of molding the fine powder to obtain a compact;
A method for producing a sintered body for a rare earth magnet, comprising the 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 longer and 50 hours or shorter to obtain a sintered body.

Images