JP7287215B2 - Manufacturing method of sintered body for rare earth magnet - Google Patents

Description

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

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 permanent 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 the substituting element M1 such as Zr and Hf, the amount of the substituting element M2 such as Ti substituting the transition metal element is reduced and saturated. A rare earth permanent magnet is disclosed that has a stabilized ThMn ₁₂ -type crystal structure while retaining its magnetization.

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 alloy is , which has a ThMn ₁₂ -type crystal structure produced by an ultra-quenching method, thereby exhibiting high magnetic properties.

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

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 included as a secondary phase in addition to the ThMn _12- type main phase. It is described that high coercive force is possible.

Patent Literature 6 describes that the addition of Cu generates a liquid phase to form a dense bulk body.

In Patent Document 7, by producing a ferromagnetic alloy having a Y-containing ThMn ₁₂ type phase as the main phase by a strip casting method, an alloy with less nonuniformity in the main phase composition and a high main phase ratio can be obtained. It is stated that

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.

Patent Literature 9 describes that a ferromagnetic alloy having high thermal stability and having a Y-containing ThMn ₁₂ type phase as a main phase can be obtained.

Patent Document 10 describes that an alloy suitable for producing anisotropic sintered magnetic powder can be obtained by adding Cu.

Patent Document 11 describes that a magnetic material having a Y-containing ThMn ₁₂ -type phase as a main phase can obtain high saturation magnetization.

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 WO2018/123988 JP 2019-44259 A JP 2019-54217 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. Also, in order to obtain high magnetic properties, a sintered body with high density 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.

The ferromagnetic alloy described in Patent Document 7, the magnetic material described in Patent Document 8, the ferromagnetic alloy described in Patent Document 9, the rare earth magnet alloy described in Patent Document 10, and the magnetic material described in Patent Document 11. Since the composition does not consider the influence of oxygen that is inevitably mixed in the manufacturing process of the sintered body, oxygen preferentially reacts with the rare earth element, the main phase decomposes, and soft magnetic such as bcc-Fe phase There is concern about the generation of phases.

Embodiments of the present disclosure provide a method for manufacturing a sintered body for a rare earth magnet with a small number of heterogeneous phases that adversely affect magnetic properties and a high density.

In an exemplary embodiment of the method for producing a sintered body for a rare earth magnet of the present disclosure,
The overall composition of the sintered body is represented by the following composition formula (1),
R(Fe _1-y Co _y ) _wz M _z Cu _α O _β (1)
R is at least one of R1 and rare earth elements other than R1, R1 is at least one selected from the group consisting of Y, Gd, Hf and Zr, M is Si, Al, Ti, V, Cr, is at least one selected from the group consisting of Nb, Mo, Ta and W, and y, z, w, α and β are respectively 0≦y≦0.4, 0.35≦z≦1.0 7≤w≤12, 0.2≤α≤1.0, 0.02≤β≤0.5, and -0.06≤1-1.45z-0.5α-0.5β≤0.02, 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 the following steps: obtaining a raw material powder; a step of heat-treating the molded body at 1160° C. or more and less than 1210° C. for 0.5 hours or more and 50 hours or less to obtain a sintered body; and performing an additional heat treatment for 50 hours or less.
In one embodiment, the composition of the sintered body contains R1, and R1 is 10 mol % or more and 70 mol % or less of the total R.
In one embodiment, the composition of the sintered body contains Sm, and Sm is 20 mol % or more and 80 mol % or less of the entire R.
In one embodiment, the composition of the sintered body contains Ti, and Ti accounts for 50 mol % or more of the total M.

According to the embodiments of the present disclosure, it is possible to provide a method for producing a sintered body for a rare earth magnet with a small number of heterogeneous phases that adversely affect magnetic properties and a high density.

[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).

R(Fe _1-y Co _y ) _wz M _z Cu _α O _β (1)
Here, R is at least one of R1 and rare earth elements other than R1, R1 is at least one selected from the group consisting of Y, Gd, Hf and Zr, M is Si, Al, Ti, V , Cr, Nb, Mo, Ta and W.
and y, z, w, α and β are 0≤y≤0.4, 0.35≤z≤1.0, 7≤w≤12, 0.2≤α≤1.0 and 0.02 respectively. ≦β≦0.5 and further satisfies the relational expression −0.06≦1−1.45z−0.5α−0.5β≦0.02.

As a result of intensive research by the present inventors, by setting the sintered body to a specific composition range as shown in the above formula (1), bcc-(Fe, Co, Ti), which adversely affects the magnetic properties, It was found that the amount of the phase and the phase of a compound having a Th ₂ Ni ₁₇ -type crystal or its similar structure (hereinafter sometimes referred to as the 2-17 phase) can be reduced. Further, the present inventors have found that a high-density sintered body can be obtained by heat-treating the molded body in a specific narrow temperature range described later when producing the sintered body of the specific composition of the present disclosure.

[Regarding the reason for limiting the composition of the sintered body, etc.]
(Type of R)
R is at least one of R1 and rare earth elements other than R1, and R1 is at least one selected from the group consisting of Y, Gd, Hf and Zr. “Rare earth elements other than R1” means rare earth elements other than Y and Gd which are rare earth elements in R1. R is an element necessary for the formation of the 1-12 phase, the Cu-containing grain boundary phase, and the oxide phase. R1 is preferably 10 mol % or more and 70 mol % or less of the entire R. Since Y, Gd, Hf and Zr play a role in stabilizing the 1-12 phase, they are preferably added to suppress the decomposition of the 1-12 phase. Furthermore, Y is more preferably 50 mol % or more of the entire R1, and most preferably R1 consists of Y. Gd and Hf are more expensive than Y. Also, when R1 is Zr, a Th ₆ Mn ₂₃ type phase and a bcc-(Fe, Co, Ti) phase associated therewith may be generated. From the viewpoint of magnetic physical properties, Sm is preferably contained in an amount of 20 mol % or more and 80 mol % or less of the entire R, and more preferably 50 mol % or more and 80 mol % or less. By including Sm in R, the 1-12 phase develops strong uniaxial anisotropy (the c-axis direction of the tetragonal crystal is the easy magnetization axis).

(Type of M)
M is at least one selected from the group consisting of Si, Al, Ti, V, Cr, Nb, Mo, Ta and W; These elements play a role in stabilizing the 1-12 phase. Preferably, Ti accounts for 50 mol % or more of the total M element, and more preferably 80 mol % or more of the total M element. Among the M elements, Ti has the effect of stabilizing the 1-12 phase even in a small amount, and can minimize the deterioration of the magnetic property values of the 1-12 phase.

(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 a decrease in 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 M)
The range of z (M content z), which indicates the atomic number ratio of the content of M to the content of R, is 0.35≦z≦1.0. When z is less than 0.35, 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 1.0, 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 M content, the better. Specifically, it is more preferable that the range of z is 0.35≦z≦0.60.

(Cu content)
The range of α (Cu content α), which indicates the atomic number ratio of the Cu content to the R content, is 0.2≦α≦1.0. If α is less than 0.2, the amount of liquid phase during heat treatment is reduced, which makes it difficult to reduce heterogeneous phases during solution treatment and densification during sintering. If α is greater than 1.0, the ratio of the R—Cu phase, which is a secondary phase, increases, 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≦1.0 in order to promote densification during sintering.

(Total amount of Fe, Co, M)
The range of w, which indicates the atomic number ratio of the total amount of Fe, Co, and M to the R content, 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.

(Content of oxygen (O))
β, which indicates the atomic number ratio of the oxygen content to the R content, 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)
z, α, and β satisfy the relational expression -0.06≤1-1.45z-0.5α-0.5β≤0.02. Since the sintered body generally uses 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 the rare earth elements in the main phase and grain boundary phase (liquid phase during sintering) during pulverization and sintering to form oxide phases. The 1-12 phase may decompose to form the bcc-(Fe, Co, Ti) phase. As a result of earnest research, the authors found out how R is distributed to each phase. From the value of z, the amount of R consumed as a 1-12 phase is 1.45z, from the value of α, the amount of R consumed as an R-Cu phase is 0.5α, and from the value of β, R oxide It is a formula that describes the R consumed as a phase as 0.5β, respectively, and defines the upper and lower limits of the actual amount of R 1 minus the amount of R calculated from z, α, and β. As 1-1.45z-0.5α-0.5β becomes smaller, it means that R necessary for forming 1-12 phase, R-Cu phase and R oxide phase is insufficient. It means that R becomes redundant. If 1-1.45z-0.5α-0.5β is less than -0.06, a large amount of bcc-(Fe, Co, Ti) phase is generated, which is not preferable. On the other hand, if it is more than 0.02, a large amount of a phase such as a 2-17 phase having a higher rare earth content than the 1-12 phase is generated, which is not preferable. In terms of the phase ratio in the sintered body, both the bcc-(Fe, Co, Ti) phase and the 2-17 phase are preferably 10% by volume or less.

[Regarding the reason for limiting the manufacturing method]
<Step A> Step of Obtaining Raw Material Powder Each element is weighed so as to obtain the above-described composition of the sintered body for a rare earth magnet to obtain a raw material powder. The raw material powder is prepared in consideration of evaporation of rare earth elements (for example, Sm) during melting and sintering. As a method for producing the raw material powder, a method of producing a raw material alloy in the form of ingots, flakes, ribbons, or the like and then pulverizing the alloy to obtain powder, or a method of directly obtaining powder by an atomizing method or the like can be employed. Known methods such as metal mold casting, centrifugal casting, strip casting, and liquid ultra-quenching can be employed as methods for producing raw material alloys in the form of ingots, flakes, ribbons, and the like. 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 heterogeneous phase contained in the alloy becomes large, it becomes difficult to eliminate the heterophase 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 and coarsening main phase grains. Depending on the composition of the alloy, the melting point of the R—Cu phase 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 to be 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 into the alloy. 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 250° C. or higher in hydrogen to absorb hydrogen, the atmosphere can be switched to a vacuum or an inert gas 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 temperatures above 700° C., decomposition of the main phase by hydrogenation-disproportionation reaction may occur. 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.

Preliminary pulverization may be performed before pulverization. For preliminary pulverization, known methods such as jaw crusher, hammer mill, and roller mill can be adopted. As a pulverization method, for example, a known method such as a jet mill, stamp mill, or ball mill can be adopted. During pre-grinding and grinding, a grinding aid may be added for efficient grinding. Known aids such as zinc stearate can be used as grinding aids. 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 powder after pulverization is _D50 (volume-based median diameter of particles when the cumulative frequency is 50%) obtained by a laser diffraction method using an airflow dispersion method. is preferably 1 μm or more and 20 μm or less. If the _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 finer the particle size of the powder and the higher the oxygen concentration in the pulverizing gas, the larger the value of the oxygen content β in the sintered body. Conversely, the coarser the particle size of the powder and the lower the oxygen concentration in the pulverization gas, the smaller the value of β. It should be noted that the pulverization step is not necessarily required when the alloy is produced by a method such as an atomizing method that can directly produce a powder. When obtaining such a powder, it may be produced from a single raw material alloy so as to obtain a desired composition of the sintered body, or it may be obtained as a mixed powder of a plurality of raw material alloys.

<Step B> Molding Step The raw material powder obtained in Step A is molded to obtain a compact. In order to orient the crystals, it is preferable to carry out molding while applying a magnetic field during molding. In addition, the molding method is a dry molding method in which dried raw material powder is inserted into the cavity of the mold and molded. A known method including a wet molding method in which molding is performed while discharging a dispersion medium can be employed.

<Step C> Sintering Step A sintered body is obtained by heat-treating the compact obtained in Step B. 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 can be employed. From the viewpoint of operating costs, etc., it is preferable to perform solid-phase sintering or liquid-phase sintering in a vacuum or an inert gas atmosphere. 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 treatment temperature is 1160° C. or higher and 1210° C. or lower. A high-density sintered body can be obtained by performing the sintering treatment in a specific narrow temperature range of 1160° C. or higher and 1210° C. or lower. If the sintering temperature is lower than 1160°C, densification will be insufficient. In addition, if the sintering temperature is higher than 1210° C., a coarse (Fe, Co, Ti) phase may occur during the sintering treatment (because it may not be a bcc structure at the sintering temperature, simply (Fe, Co, Ti) phase) and 2-17 phase may be generated, and the heterogeneous phase may not be sufficiently reduced even in the subsequent additional heat treatment process. The sintering treatment temperature is more preferably 1180° C. or higher and 1210° C. or lower. The sintering treatment time is 0.5 hours or more and 50 hours or less. If the sintering treatment time is less than 0.5 hours, 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. Moreover, when the compact is wet, it is better to deoil it at a temperature at which oil evaporates before reaching the sintering temperature.

<Step D> Additional heat treatment By additionally heat-treating the sintered body obtained in step D, a sintered body with less heterogeneous phases is obtained. In the sintering temperature range in step C, not only the 1-12 phase but also the 2-17 phase and (Fe, Co, Ti) phase are stable, so the number of different phases increases during sintering. Therefore, an additional heat treatment is performed in a temperature range in which the 1-12 phase is stable to promote the reaction of the 2-17 phase + (Fe, Co, Ti) phase → 1-12 phase and reduce the heterogeneous phase. The heat treatment temperature is 900° C. or higher and 1150° C. or lower. If the heat treatment temperature is less than 900° C., the atoms are not sufficiently diffused, and the reaction of 2-17 phase+(Fe, Co, Ti) phase→1-12 phase is difficult to proceed. Further, if the heat treatment temperature exceeds 1150° C., the 2-17 phase and the α-Fe phase are also in a stable region, so the reduction of heterogeneous phases is insufficient, which is not suitable. The heat treatment temperature is more preferably 950° C. or higher and 1120° C. or lower. The heat treatment time is 0.5 hours or more and 50 hours or less. If the heat treatment time is less than 0.5 hours, there is a possibility that the reduction of heterogeneous phases will not proceed sufficiently. Moreover, if the heat treatment time exceeds 50 hours, the lead time becomes long, which is not preferable in terms of operation.
The sintered body after the additional heat treatment may be additionally subjected to heat treatment or diffusion treatment of a specific element for the purpose of improving the coercive force.

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

Experimental example 1
First, each element was weighed so as to have the compositions of A1, A2 and A3 in Table 1 as main phase material alloys, and strip casting was performed. Specifically, raw material metals of Y, Sm, Fe, Co, Ti, and Cu with a purity of 99.9% or more (Si is contained as an unavoidable impurity during the manufacturing process) are added to the rare earth element at the time of melting. Evaporation was taken into account and weighing was carried out so that the obtained alloy composition would be the target value. 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.

Next, each element was weighed so that the composition of B1 in Table 1 was obtained as a subphase alloy, and the alloy was produced by a super-quenching method. Specifically, raw material metals of Y, Sm, Fe, Co, Ti, and Cu with a purity of 99.9% or more (Si is contained as an unavoidable impurity during the manufacturing process) are added to the rare earth element at the time of melting. Evaporation was taken into account and weighing was carried out so that the obtained alloy composition would be the target value. After sufficiently melting these raw metals in the tapping pipe of a liquid ultra-quenching device (melt spinning device) to form the core of the alloy, the molten metal is poured onto a Cu roll rotating at a roll peripheral speed of 20 m/s. I took a bath.

A portion of the main phase alloys A1 to A3 and the sub phase alloy B1 thus prepared were each pulverized in a mortar and classified using a 425 μm mesh and a 75 μm mesh. Using pulverized powder with a particle size of 75 to 425 μm, component analysis of Y, Sm, Fe, Co, Ti, Cu, and Si was performed by ICP (inductively coupled plasma) emission spectrometry. Table 1 shows the composition of each alloy.

The main phase system alloys A1 to A3 thus produced were subjected to alloy heat treatment. Specifically, 500 g of each 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 finishing the alloy heat treatment, 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 main phase alloys A1 to A3 and the sub phase alloy B1 after alloy heat treatment obtained in the above steps were placed in a molybdenum container, and the container was inserted into a tubular heat treatment furnace. After replacing the atmosphere in the furnace with H ₂ , heat treatment was performed at 350° C. for 1.5 hours in an atmosphere in which 2 L/min of H ₂ was allowed to flow. After the heat treatment was finished, the atmosphere in the furnace was replaced with Ar, and then the heat treatment furnace was opened to cool the alloy.

The hydrogen-treated main phase alloy and sub phase alloy obtained in the above steps were pulverized using a mortar in a glove box in an Ar stream atmosphere. The pulverized powder was sieved through a 1 mm mesh, and the powder that passed through the mesh was recovered. The collected main phase alloy powder and sub phase alloy powder and zinc stearate (grinding aid for the next step) were mixed in a rocking mixer for 20 minutes.

The mixed powder obtained in the above step was pulverized using an air 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 pulverization pressure of 7.0 MPa to obtain powder. _D50 of the powder at this time was 5 μm.

The powder obtained in the above step was mixed with oil to form a slurry, which was compacted in a magnetic field to obtain a compact. A so-called perpendicular magnetic field forming device (transverse magnetic field forming device) was used as the forming device, in which the direction of magnetic field application and the direction of pressure are orthogonal. The molded body obtained was covered with a niobium foil.

The molded body obtained in the above process was placed in a molybdenum container, deoiled in a heat treatment furnace in a vacuum atmosphere at 200 ° C. for 5 hours, then filled with Ar, and the sintering treatment temperature shown in Table 2. 20 hours of sintering process.

A portion of the sintered body obtained in the above step was 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, an additional heat treatment was performed at 1100° C. for 20 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.

A portion of the sintered body obtained in the above step was pulverized in a mortar and classified using a 425 μm mesh and a 75 μm mesh. Using pulverized powder with a particle size of 75 to 425 μm, the components of Y, Sm, Fe, Co, Ti, Cu, and Si are analyzed by ICP emission spectrometry, and the carbon content is analyzed by combustion and infrared absorption. rice field. In addition, using pulverized powder having a particle size of 425 μm or more, the oxygen content and nitrogen content were analyzed by the inert gas fusion/heat conduction method. From the analysis results, the values of y, z, w, α, β, and 1-1.45z-0.5α-0.5β of each sintered body were determined. Also, the density of the sintered body was obtained by the Archimedes method using ion-exchanged water. It was confirmed by a powder X-ray diffraction method using pulverized powder having a particle size of 75 μm or less that a phase having a ThMn type ₁₂ crystal structure was the main phase.

The sintered body was cut with a peripheral blade cutting machine. The cut sintered body was embedded in resin, polished, and the cross section of the sintered body was observed with a scanning electron microscope (SEM). A JCM-6000Plus NeoScope (registered trademark) manufactured by JEOL Ltd. was used as the SEM, and backscattered electron images were taken at an acceleration voltage of 15 kV. The captured backscattered electron image was analyzed using image processing software. Based on the contrast of each phase, the cross-sectional area ratios of the bcc-(Fe, Co, Ti) phase and the 2-17 phase were obtained in a region of 400 μm×600 μm.

Types of main phase alloy and sub phase alloy used in the produced sintered body, mixing weight ratio of the sub phase alloy when the weight of the main phase alloy is 1, composition of the sintered body, y, z , w, α, β, 1-1.45z-0.5α-0.5β values, sintering temperature, presence or absence of additional heat treatment, sintered body density, bcc-(Fe, Co, Ti) phase ratio and Table 2 shows the 2-17 phase ratio. In addition, No. Samples Nos. 21 to 25 used only the main phase system alloy and did not mix the sub phase system alloy.

No. 1, 6, 11, 16 and 21 are experimental examples in which the sintering temperature is 1150°C. No. Compared with samples other than 1, 6, 11, 16 and 21, the sintered body density was remarkably low.

No. Samples 2, 3, 7, 8, 12, 13, 17, 18, 22, and 23 were sintered at 1200°C or 1220°C without additional heat treatment. Although the density of the sintered body was high, either or both of the bcc-(Fe, Co, Ti) phase ratio and the 2-17 phase ratio exceeded 10%, resulting in an extremely large number of different phases.

No. Nos. 4, 9, 19 and 24 are samples which were subjected to additional heat treatment at 1100°C after sintering at 1200°C. The density of the sintered body was high, and both the bcc-(Fe, Co, Ti) phase ratio and the 2-17 phase ratio could be suppressed to 10% or less.

No. 14 is No. Similar to 4, 9, 19, and 24, after sintering at 1200°C, additional heat treatment was performed at 1100°C. Yes, samples higher than 0.02. Although the density of the sintered body is high, the 2-17 phase ratio exceeds 10%, resulting in an extremely large number of different phases.

No. 5, 10, 15, 20, and 25 are samples which were subjected to additional heat treatment at 1100°C after sintering at 1220°C. Although the density of the sintered body was high, either or both of the bcc-(Fe, Co, Ti) phase ratio and the 2-17 phase ratio exceeded 10%, resulting in an extremely large number of different phases.

The sintered body for rare earth magnets of the present disclosure is suitably used in various technical fields where permanent magnets that exceed the magnetic properties of Nd--Fe--B magnets are required, especially motors and actuators, etc., and various industrial applications. have a use.

Claims (4)

The overall composition of the sintered body is represented by the following composition formula (1),
R(Fe _1-y Co _y ) _wz M _z Cu _α O _β (1)
R is at least one of R1 and rare earth elements other than R1, R1 is at least one selected from the group consisting of Y, Gd, Hf and Zr,
M is at least one selected from the group consisting of Si, Al, Ti, V, Cr, Nb, Mo, Ta and W;
y, z, w, α and β are each
0≤y≤0.4,
0.35≦z≦1.0,
7≦w≦12,
0.2≤α≤1.0,
0.02≦β≦0.5, and −0.06≦1−1.45z−0.5α−0.5β≦0.02,
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 obtaining raw material powder;
a step of molding the raw material powder to obtain a molded body;
a step of heat-treating the molded body at 1160° C. or higher and less than 1210° C. for 0.5 hours or more and 50 hours or less to obtain a sintered body;
a step of subjecting the sintered body to additional heat treatment at 900° C. or more and less than 1150° C. for 0.5 hours or more and 50 hours or less;
A method for producing a sintered body for rare earth magnets. 2. The method for producing a sintered body for a rare earth magnet according to claim 1, wherein the composition of said sintered body contains R1, and R1 is 10 mol % or more and 70 mol % or less of the total R. 3. The method for producing a sintered body for a rare earth magnet according to claim 1, wherein the composition of said sintered body contains Sm, and Sm is 20 mol % or more and 80 mol % or less of the whole R. 4. The method for producing a sintered body for a rare earth magnet according to any one of claims 1 to 3, wherein the composition of the sintered body contains Ti, and Ti accounts for 50 mol% or more of the total M.