JP2020129679A - R-Fe-B system sintered magnet - Google Patents

Abstract

SOLUTION: (R', HR) 2 (Fe, (Co)) 14B (R'is one or more elements selected from rare earth elements excluding Dy, Tb and Ho, including Y, and is Nd. HR is a main phase containing an HR rich phase represented by (one or more elements selected from Dy, Tb and Ho) and 25 to 35 atomic% (R', HR), 2 to 8 atoms. % M1 (M1 is one or more selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. Element), grain boundary phase containing (R', HR) -Fe (Co) -M1 phase having a composition of Co of 8 atomic% or less and Fe of the balance in one or both forms of an amorphous phase and a microcrystalline phase. An R-Fe-B based sintered magnet having a high HR content in the HR rich phase and higher than the HR content in the central part of the main phase. [Effect] The R-Fe-B-based sintered magnet of the present invention gives a high coercive force even if the contents of Dy, Tb and Ho are small. [Selection diagram] Fig. 1

Description

The present invention relates to an R-Fe-B based sintered magnet having a high coercive force.

Nd-Fe-B system sintered magnets (hereinafter sometimes referred to as Nd magnets) have been expanding their application range and production year by year as functional materials indispensable for energy saving and high functionality. For example, since it is expected to be used in a high temperature environment for automobile applications, for example, an Nd magnet incorporated in a drive motor of a hybrid vehicle or an electric vehicle or an electric power steering motor has a high residual magnetic flux density. At the same time, high coercive force is required. On the other hand, the coercive force of the Nd magnet tends to be remarkably reduced at high temperatures, and in order to secure the coercive force at the operating temperature, it is necessary to sufficiently increase the coercive force at room temperature in advance.

As a method of increasing the coercive force of the Nd magnet, it is effective to substitute a part of Nd of the main phase Nd ₂ Fe ₁₄ B compound with Dy or Tb, but these elements have a small resource reserve. Rather, there is a risk that prices will be unstable and fluctuations will be large because the commercially viable production areas will be limited and the stable supply will be affected by geopolitical factors. From such a background, in order to obtain a large market for R-Fe-B based magnets that can be used at high temperatures, a new method of increasing the coercive force after suppressing the addition amount of Dy or Tb as much as possible or R -Fe-B magnet composition needs to be developed. From this point of view, various methods have been conventionally proposed.

For example, in Japanese Patent No. 3997413 (Patent Document 1), R of 12 to 17% in atomic percentage (R is at least two or more of rare earth elements including Y, and Nd and Pr are essential), 0.1 to 3% Si, 5 to 5.9% B, 10% or less Co and the balance Fe (provided that Fe is Al, Ti, V, Cr, Mn, Ni with a substitution amount of 3 at% or less). , Cu, Zn, Ga, Ge, Zr, Nb, Mo, In, Sn, Sb, Hf, Ta, W, Pt, Au, Hg, Pb, Bi, even if it is substituted with one or more elements. In a sintered R—Fe—B based magnet having a coercive force of at least 10 kOe or more and having a composition of R ₂ (Fe, (Co), Si) ₁₄ B as a main phase, The R-Fe(Co)-Si grain boundary phase, which does not include a rich phase and is composed of 25 to 35% R in atomic percentage, 2 to 8% Si, 8% or less Co, and the balance Fe, is at least in volume ratio. An R-Fe-B based sintered magnet having 1% or more of the entire magnet is disclosed. This sintered magnet is cooled at a rate of 0.1 to 5° C./minute during at least 700 to 500° C. in a cooling step during the sintering or the heat treatment after the sintering of its manufacture. Alternatively, the R-Fe(Co)-Si grain boundary phase is formed in the structure by maintaining a constant temperature for at least 30 minutes or more and cooling in multiple stages during cooling.

Japanese Patent Publication No. 2003-510467 (Patent Document 2) discloses an Nd-Fe-B alloy having a low boron content, a sintered magnet using this alloy, and a method for manufacturing the same, and a sintered magnet is manufactured from this alloy. As a method to do so, when the raw material is cooled to 300° C. or lower after sintering, the average cooling rate up to 800° C. is cooled at ΔT ₁ /Δt ₁ <5 K/min.

Japanese Patent No. 5572673 (Patent Document 3), R-T-B magnet and a R ₂ Fe ₁₄ B main phase and a grain boundary phase is described. Part of this grain boundary phase is an R-rich phase containing more R than the main phase, and the other grain boundary phases are transition metal-rich phases having a lower rare earth element concentration and a higher transition metal element concentration than the main phase. is there. It is described that the R-T-B rare earth sintered magnet is manufactured by performing sintering at 800°C to 1200°C and then performing heat treatment at 400°C to 800°C.

In Japanese Patent Laid-Open No. 2014-132628 (Patent Document 4), the grain boundary phase is an R-rich phase in which the total atomic concentration of rare earth elements is 70 atomic% or more, and the total atomic concentration of rare earth elements is 25 to 35 atomic %. And an R-T-B rare earth sintered magnet having an area ratio of the transition metal-rich phase in the grain boundary phase of 40% or more, which includes a transition metal-rich phase that is ferromagnetic, and describes the magnetic alloy pressure. After performing the step of sintering the powder compact at 800° C. to 1,200° C. and the first heat treatment step at 650° C. to 900° C., cooling to 200° C. or lower, and further performing the second heat treatment step at 450 It is described to be manufactured by a plurality of heat treatment steps performed at a temperature of from 600°C to 600°C.

Japanese Unexamined Patent Application Publication No. 2014-146788 (Patent Document 5) discloses an R-T-B rare earth sintered magnet including a main phase made of R ₂ Fe ₁₄ B and a grain boundary phase containing more R than the main phase. , The easy axis of magnetization of the R ₂ Fe ₁₄ B main phase is parallel to the c-axis, and the crystal grain shape of the R ₂ Fe ₁₄ B main phase is an elliptical shape extending in the direction orthogonal to the c-axis direction, and the grain boundary phase , An R-T-B rare earth sintered magnet containing an R-rich phase in which the total atomic concentration of rare earth elements is 70 atomic% or more, and a transition metal-rich phase in which the total atomic concentration of rare earth elements is 25 to 35 atomic %. Is listed. Further, in the production, it is described that sintering is performed at 800° C. to 1,200° C., and after sintering, heat treatment is performed at 400° C. to 800° C. in an argon atmosphere.

Japanese Unexamined Patent Application Publication No. 2014-209546 (Patent Document 6) includes a R ₂ T ₁₄ B main phase and a two-particle grain boundary phase between two adjacent R ₂ T ₁₄ B main phase crystal grains. A rare-earth magnet is disclosed which has a thickness of the two-grain grain boundary phase of 5 nm or more and 500 nm or less and which has a magnetic property different from that of a ferromagnetic material. This rare earth magnet is formed from a compound that contains T element as a two-grain grain boundary phase but does not become ferromagnetic. Therefore, this phase contains a transition metal element and is composed of Al, Ge, Si, It contains M elements such as Sn and Ga. Further, by adding Cu to the rare earth magnet, a crystal phase having a La ₆ Co ₁₁ Ga ₃ type crystal structure can be uniformly and widely formed as a two-grain grain boundary phase, and at the same time, a La ₆ Co ₁₁ Ga ₃ type two-grain grain boundary phase can be formed. A thin R-Cu layer can be formed at the interface between the R ₂ T ₁₄ B and the crystal grains of the R ₂ T ₁₄ B main phase, thereby passivating the interface of the main phase, suppressing the occurrence of strain due to lattice mismatch, and suppressing the reverse magnetic domain. It is described that it can be suppressed to become the nucleus of generation of. Then, in the production thereof, post-sintering heat treatment is performed at 500° C. to 900° C., and cooling is performed at a cooling rate of 100° C./min or more, particularly 300° C./min or more.

In International Publication No. 2014/157448 (Patent Document 7) and International Publication No. 2014/157451 (Patent Document 8), an Nd ₂ Fe ₁₄ B type compound is used as a main phase, and it is surrounded by two main phases and has a thickness. An RTB-based sintered magnet having a two-grain grain boundary of 5 to 30 nm and a grain boundary triple point surrounded by three or more main phases is disclosed.

Japanese Patent No. 3997413 Special table 2003-510467 gazette Japanese Patent No. 5572673 JP, 2014-132628, A JP, 2014-146788, A JP, 2014-209546, A International Publication No. 2014/157448 International Publication No. 2014/157451

From the above-mentioned circumstances, there is a demand for an R—Fe—B based sintered magnet that exhibits a high coercive force even if the contents of Dy, Tb, and Ho are small.

The present invention has been made in view of the above circumstances, and an object thereof is to provide a novel R-Fe-B based sintered magnet having a high coercive force.

As a result of intensive studies to solve the above problems, the present inventors have found that 12 to 17 atom% of R (R is one or more elements selected from rare earth elements including Y, and Nd is 0.1 to 3 atomic% of M ₁ (M ₁ is Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au). , Hg, Pb and Bi), 0.05 to 0.5 atomic% of M ₂ (M ₂ is Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W). One or more elements selected), (4.8+2×m to 5.9+2×m) atomic% (m is the content (atomic %) of the element represented by M ₂ ) B, 10 atomic% or less R ₂ (Fe, (Co)) ₁₄ B having a composition of Co, C of 0.5 atomic% or less, O of 1.5 atomic% or less, N of 0.5 atomic% or less, and the balance of Fe. It is an R-Fe-B type|system|group sintered magnet which uses an intermetallic compound as a main phase, is formed between the main phase and the crystal grain of a main phase, and is 25-35 atomic% of (R', HR), 2-. (R′,HR)—Fe(Co)—M ₁ phase (R′ contains Y, Dy, Tb and Ho) having a composition of 8 atomic% M ₁ , 8 atomic% or less Co, and the balance Fe. , Which is one or more elements selected from rare earth elements other than Nd, and Nd is essential, and HR is one or more elements selected from Dy, Tb, and Ho), an amorphous phase, and a fine particle having a particle size of 10 nm or less. An HR-rich phase represented by (R′,HR) ₂ (Fe,(Co)) ₁₄ B on the surface thereof, which has a grain boundary phase contained in one or both of the crystal phases. The R-Fe-B system sintered magnet containing HR in the HR rich phase has a higher HR content than the HR content in the central portion of the main phase is an R-Fe-B system sintered body having a high coercive force. It was found to be a magnet.

Then, in such an R—Fe—B system sintered magnet, a step of preparing a fine alloy powder having a predetermined composition, a step of compacting the fine alloy powder in a magnetic field to obtain a compact, a compact 900 ˜1,250° C. sintering to obtain a sintered body, cooling the sintered body to a temperature of 400° C. or lower, HR (HR is one or more selected from Dy, Tb and Ho) A metal, a compound or an intermetallic compound containing the element) is heated on the surface of the sintered body at a temperature in the range of more than 950° C. and not more than 1,100° C. It was found that it can be produced by a method including a high temperature heat treatment step of diffusing and cooling to 400° C. or lower, and a low temperature heat treatment step of cooling to 300° C. or lower by heating at a temperature in the range of 400 to 600° C. after high temperature heat treatment, Invented.

Therefore, the present invention provides the following R-Fe-B system sintered magnet.
[1] 12 to 17 atomic% of R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic% of M ₁ ( M ₁ is one or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi), 0.05 to 0.5 atomic% of M ₂ (M ₂ is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W), (4.8+2×m 5.9+2×m) atomic% (m is the content (atomic %) of the element represented by M ₂ ) B, 10 atomic% or less Co, 0.5 atomic% or less C, 1.5 atomic% R-Fe-B based sintering having the following composition of O, N of 0.5 atomic% or less, and the balance of Fe, and having R ₂ (Fe, (Co)) ₁₄ B intermetallic compound as a main phase. A magnet,
Formed between the main phase and the crystal grains of the main phase, 25 to 35 atomic% of (R′, HR), 2 to 8 atomic% of M ₁ , 8 atomic% or less of Co, and the balance of Fe (R′,HR)-Fe(Co)-M ₁ phase having a composition (R′ includes Y, and is one or more elements selected from rare earth elements other than Dy, Tb and Ho, and Nd. HR has a grain boundary phase containing one or both of an amorphous phase and a microcrystalline phase having a grain size of 10 nm or less, and HR has one or more elements selected from Dy, Tb and Ho.
(R ', HR) M ₁ in -Fe (Co) -M ₁ phase,
0.5 to 50 atomic% is Si, and the balance is selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements,
1.0 to 80 atomic% is Ga, and the rest is Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements, or 0.5 to 50 atomic% of Al and the balance of Si, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg. , One or more elements selected from Pb and Bi,
The main phase includes an HR rich phase represented by (R′,HR) ₂ (Fe,(Co)) ₁₄ B on the surface thereof, and the content of HR in the HR rich phase is An R-Fe-B based sintered magnet having a higher HR content in the center of the phase.
[2] The R-Fe-B based sintered magnet according to [1], wherein the HR-rich phase is unevenly formed on the surface of the main phase.
[3] The content of Nd in the HR-rich phase is 0.8 times or less the content of Nd in the central portion of the main phase, R- according to [1] or [2]. Fe-B system sintered magnet.
[4] The area of the HR rich phase evaluated in a cross section within 200 μm from the surface of the R—Fe—B system sintered magnet is 2% or more with respect to the area of the entire main phase [ 1] to the R-Fe-B system sintered magnet according to any one of [3].
[5] Any of [1] to [4], wherein the (R′,HR)-Fe(Co)-M ₁ phase is present in the structure of the magnet in a volume ratio of 1% or more and 20% or less. The R-Fe-B type|system|group sintered magnet of crab.
[6] above (R ', HR) M ₁ in -Fe (Co) -M ₁ phase,
0.5 to 50 atomic% is Si, and the balance is selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements, or 1.0 to 80 atomic% Ga, the balance Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg , Pb and Bi are one or more elements, and the R-Fe-B system sintered magnet according to any one of [1] to [5] is characterized.
[7] In any one of [1] to [6], the residual magnetic flux density Br is 13.2 kG or more at 23° C., and the coercive force Hcj is 26.2 kOe or more at 23° C. The described R-Fe-B based sintered magnet.
The present invention also relates to the following R-Fe-B based sintered magnet.
[1] 12 to 17 atomic% of R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic% of M ₁ ( M ₁ is one or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi), 0.05 to 0.5 atomic% of M ₂ (M ₂ is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W), (4.8+2×m 5.9+2×m) atomic% (m is the content (atomic %) of the element represented by M ₂ ) B, 10 atomic% or less Co, 0.5 atomic% or less C, 1.5 atomic% R-Fe-B based sintering having the following composition of O, N of 0.5 atomic% or less, and the balance of Fe, and having R ₂ (Fe, (Co)) ₁₄ B intermetallic compound as a main phase. A magnet,
Formed between the main phase and the crystal grains of the main phase, 25 to 35 atomic% of (R′, HR), 2 to 8 atomic% of M ₁ , 8 atomic% or less of Co, and the balance of Fe (R′,HR)-Fe(Co)-M ₁ phase having a composition (R′ includes Y, and is one or more elements selected from rare earth elements other than Dy, Tb and Ho, and Nd. HR has a grain boundary phase containing one or both of an amorphous phase and a microcrystalline phase having a particle size of 10 nm or less), and HR has one or more elements selected from Dy, Tb and Ho.
The main phase contains an HR rich phase represented by (R′,HR) ₂ (Fe,(Co)) ₁₄ B in its surface portion,
The R-Fe-B based sintered magnet, wherein the content of HR in the HR-rich phase is higher than the content of HR in the central portion of the main phase.
[2] The R-Fe-B based sintered magnet according to [1], wherein the HR-rich phase is formed unevenly on the surface of the main phase.
[3] The content of Nd in the HR-rich phase is 0.8 times or less the content of Nd in the central portion of the main phase, R- according to [1] or [2]. Fe-B system sintered magnet.
[4] The area of the HR rich phase evaluated in a cross section within 200 μm from the surface of the R—Fe—B system sintered magnet is 2% or more with respect to the area of the entire main phase [1. ] The R-Fe-B type|system|group sintered magnet in any one of [3].
[5] 12 to 17 atom% of R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atom% of M ₁ ( M ₁ is one or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi), 0.05 to 0.5 atomic% of M ₂ (M ₂ is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W), (4.8+2×m-5) .9+2×m) atomic% (m is the content (atomic %) of the element represented by M ₂ ) B, 10 atomic% or less Co, 0.5 atomic% or less C, 1.5 atomic% or less Of O, 0.5 atomic% or less of N, and the balance of Fe having a composition of:
A step of compacting the alloy fine powder in a magnetic field to obtain a compact,
A step of sintering the molded body at a temperature in the range of 900 to 1,250° C. to obtain a sintered body,
Cooling the sintered body to a temperature of 400° C. or lower,
A metal, compound or intermetallic compound containing HR (HR is one or more elements selected from Dy, Tb and Ho) is arranged on the surface of the above-mentioned sintered body, and the temperature is higher than 950°C and lower than 1,100°C. After the high temperature heat treatment step of heating at a temperature in the range to diffuse the HR into the sintered body at the grain boundaries and cooling to 400° C. or less, and after the high temperature heat treatment, the temperature is in the range of 400 to 600° C. An R-Fe-B based sintered magnet obtained by a manufacturing method including a low temperature heat treatment step of cooling to the following.

The R-Fe-B system sintered magnet of the present invention gives a high coercive force even if the contents of Dy, Tb and Ho are small.

3 is an image obtained by observing an element distribution within 200 μm from a diffusion surface of the sintered magnet manufactured in Example 2 with an electron probe microanalyzer (EPMA), (A) showing Nd distribution, and (B) showing Tb distribution. The distribution is shown. 3 is an image obtained by observing the element distribution within 200 μm from the diffusion surface of the sintered magnet produced in Comparative Example 2 with an electron probe microanalyzer (EPMA), (A) showing the distribution of Nd, and (B) showing the distribution of Tb. The distribution is shown.

Hereinafter, the present invention will be described in more detail.
First, to describe the magnet composition of the present invention, R-Fe-B based sintered magnet of the present invention, 12 to 17 atomic% of R, 0.1 to 3 atomic% of M _1, 0.05 to 0. 5 atomic% of M ₂ , B of (4.8+2×m to 5.9+2×m) atomic% (m is the content rate (atomic %) of the element represented by M ₂ ) and 10 atomic% or less of Co, It has a composition of C (carbon) of 0.5 atomic% or less, O (oxygen) of 1.5 atomic% or less, N (nitrogen) of 0.5 atomic% or less, and the balance Fe, and contains inevitable impurities. You can leave.

R is one or more elements selected from rare earth elements including Y, and Nd is essential. As the rare earth element other than Nd, Pr, La, Ce, Gd, Dy, Tb and Ho are preferable, and Pr, Dy, Tb and Ho are particularly preferable, and Pr is particularly preferable. The content of R is 12 to 17 atomic %, preferably 13 atomic% or more, and 16 atomic% or less. When the content of R is less than 12 atom %, the coercive force of the magnet is extremely reduced, and when it exceeds 17 atom %, the residual magnetic flux density Br is reduced. In R, the ratio of Nd, which is an essential component, is preferably 60% by atom or more, and particularly preferably 70% by atom or more of the entire R. When one or more elements selected from Pr, La, Ce and Gd are contained as rare earth elements other than Nd, the ratio of Nd to one or more elements selected from Pr, La, Ce and Gd ( The atomic ratio of Nd/(one or more elements selected from Pr, La, Ce and Gd) is preferably 75/25 or more and 85/15 or less. In particular, when Pr is used as the rare earth element other than Nd, it is possible to use didymium which is a mixture of Nd and Pr, and in that case, the ratio Nd/Pr (atomic ratio) of Nd and Pr is, for example, 77/ It can be 23 or more and 83/17 or less.

The content of one or more elements selected from Dy, Tb and Ho is preferably 20 atom% or less of the total R, more preferably 10 atom% or less, further preferably as the total of Dy, Tb and Ho. Is 5 atomic% or less, particularly preferably 3 atomic% or less, and preferably 0.06 atomic% or more. On the other hand, the total content of Dy, Tb and Ho with respect to the composition of the entire magnet is preferably 3 atom% or less, more preferably 1.5 atom% or less, further preferably 1 atom% or less, particularly preferably It is 0.4 atom% or less, preferably 0.01 atom% or more. When one or more elements selected from Dy, Tb and Ho are diffused by grain boundary diffusion, the amount of diffusion is preferably 0.7 atom% or less, particularly 0.4 atom% or less, and 0.05 atom% or more. Is more preferable.

M ₁ is one or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. Composed. M ₁ is an element necessary for forming a (R′,HR)-Fe(Co)-M ₁ phase described later, and by adding M ₁ at a predetermined content, (R′,HR)-Fe The (Co)-M ₁ phase can be stably formed. The content of M ₁ is 0.1 to 3 atom %, preferably 0.5 atom% or more, and preferably 2.5 atom% or less. When the content of M ₁ is less than 0.1 atomic %, the coercive force is sufficiently improved because the abundance ratio of the (R′,HR)—Fe(Co)—M ₁ phase in the grain boundary phase is too low. However, if it exceeds 3 atomic %, the squareness of the magnet is deteriorated and the residual magnetic flux density Br is lowered, which is not preferable.

M ₂ is composed of at least one element selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W. M ₂ is added as an element for stably forming boride for the purpose of suppressing abnormal grain growth during sintering. The content rate of M ₂ is 0.05 to 0.5 atom %. By adding M ₂ , it becomes possible to sinter at a relatively high temperature during manufacturing, which leads to improvement of squareness and improvement of magnetic properties.

The content rate of B (boron) is (4.8+2×m) to (5.9+2×m) atomic% (m is the content rate of the element represented by M ₂ (atomic %), the same applies hereinafter), It is preferably (4.9+2×m) atom% or more, and preferably (5.7+2×m) atom% or less. In other words, since the content of the M ₂ element in the composition of the magnet of the present invention is from 0.05 to 0.5 atomic%, the content of B by the content of the M ₂ element identified in the above range Although the ranges will be different, the B content is 4.9 to 6.9 atomic %, preferably 5.0 atomic% or more, and 6.7 atomic% or less. preferable. The upper limit of the B content is an important factor. When the content ratio of B exceeds (5.9+2×m) atom %, the (R′,HR)-Fe(Co)-M ₁ phase described later is not formed at the grain boundary, and the R _1.1 Fe ₄ B ₄ compound is formed. Phase or (R′,HR) _1.1 Fe ₄ B ₄ compound phase, so-called B-rich phase is formed. When this B-rich phase exists in the magnet, the coercive force of the magnet does not increase sufficiently. On the other hand, when the content ratio of B is less than (4.8+2×m) atom %, the volume ratio of the main phase is lowered and the magnetic properties are lowered.

Co may or may not be contained, but for the purpose of improving the Curie temperature and the corrosion resistance, Fe can be replaced by Co. When Co is contained, the Co content rate is Is preferably 10 atomic% or less, and particularly preferably 5 atomic% or less. If the Co content exceeds 10 atomic %, the coercive force may be significantly reduced. The content ratio of Co is more preferably 10 atom% or less, and particularly preferably 5 atom% or less with respect to the total of Fe and Co. In the present invention, “Fe, (Co)” or “Fe(Co)” is used as a notation that means that both the case of containing Co and the case of not containing Co are included.

The carbon, oxygen and nitrogen contents are preferably lower and more preferably not contained, but in the manufacturing process, contamination cannot be completely avoided. Regarding the content of these elements, the content of C (carbon) is 0.5 atom% or less, particularly 0.4 atom% or less, and the content of O (oxygen) is 1.5 atom% or less, particularly 1.2. The content of N (nitrogen) is 0.5 atomic% or less, and especially 0.3 atomic% or less is acceptable. The content of Fe, which is the balance with respect to the composition of the entire magnet, is preferably 70 atom% or more, particularly 75 atom% or more, and 85 atom% or less, particularly 80 atom% or less.

In addition to these elements, the inclusion of elements such as H, F, Mg, P, S, Cl, and Ca as unavoidable impurities as the total of the unavoidable impurities with respect to the total of the constituent elements of the magnet and the unavoidable impurities described above. Although it is allowed up to 0.1% by mass or less, it is preferable that the content of these unavoidable impurities is small.

The average diameter of the crystal grains of the R—Fe—B system sintered magnet of the present invention is preferably 6 μm or less, particularly 5.5 μm or less, especially 5 μm or less, and more preferably 1.5 μm or more, especially 2 μm or more. preferable. The average diameter of the crystal grains can be controlled by adjusting the average particle diameter of the fine alloy powder during fine pulverization. The average diameter of crystal grains can be measured, for example, by the following procedure. First, after polishing the cross section of the sintered magnet to a mirror surface, it is immersed in an etching solution such as a villera solution (for example, a mixture solution of glycerin:nitric acid:hydrochloric acid=3:1:2) for grain boundaries. The cross section in which the phases are selectively etched is observed with a laser microscope. Next, based on the obtained observation image, the cross-sectional area of each particle is measured by image analysis, and the diameter as an equivalent circle is calculated. Then, the average diameter is obtained based on the data of the area fraction occupied by each particle size. The average diameter may be, for example, the average of a total of about 2,000 particles in 20 different images.

The residual magnetic flux density Br of the R-Fe-B system sintered magnet of the present invention is 11 kG (1.1 T) or more, particularly 11.5 kG (1.15 T) or more, especially 12 kG (1. It is preferably 2T) or more. On the other hand, the coercive force Hcj of the R—Fe—B system sintered magnet of the present invention is 10 kOe (796 kA/m) or more at room temperature (about 23° C.), particularly 14 kOe (1,114 kA/m) or more, and particularly 16 kOe (1 , 274 kA/m) or more.

The structure of the magnet of the present invention has a main phase (crystal grains) and a grain boundary phase. The main phase forming the crystal grains includes a phase of R ₂ (Fe,(Co)) ₁₄ B intermetallic compound. Here, a case not containing Co can be expressed as R ₂ Fe ₁₄ B, and a case containing Co can be expressed as R ₂ (Fe,Co) ₁₄ B.

On the other hand, the HR-rich phase contained in the main phase contains (R′,HR) ₂ (Fe,(Co)) ₁₄ B (R′ includes Y and is selected from rare earth elements other than Dy, Tb and Ho 1 A phase represented by one kind or two or more kinds of elements, Nd being essential (hereinafter the same), and HR being one or more kinds of elements selected from Dy, Tb and Ho (the same below) is included. Here, a case not containing Co can be expressed as (R′,HR) ₂ Fe ₁₄ B, and a case containing Co can be expressed as (R′,HR) ₂ (Fe,Co) ₁₄ B. The HR-rich phase is an intermetallic compound phase having a higher HR content than the HR content in the central portion of the main phase. As the rare earth element other than Nd in R′, Pr, La, Ce and Gd are preferable, and Pr is particularly preferable. The HR rich phase is formed on the surface of the main phase.

The HR rich phase is preferably formed nonuniformly on the surface of the main phase. The HR-rich phase may be formed on the entire surface of the main phase, for example, may be formed so as to cover the entire part of the main phase other than the HR-rich phase (ie, the inside). In that case, the thickness of the HR rich phase is not uniform and preferably has the thickest part and the thinnest part. In this case, the ratio of the thickest part to the thinnest part is preferably 1.5 times or more, particularly 2 times or more, and more preferably 3 times or more.

Further, the HR rich phase may be formed only on a part of the surface portion of the main phase, for example, may be formed so as to cover only a part of the portion of the main phase other than the HR rich phase. .. In this case, the thickness of the thickest part of the HR-rich phase is 0.5% or more, particularly 1% or more, especially 2% or more, and 40% or less, especially 30% or less of the grain size of the crystal grains of the main phase, It is particularly preferably 25% or less.

The thickness of the thinnest portion of the HR-rich phase is preferably 0.01 μm or more, particularly preferably 0.02 μm or more, and the thickness of the thickest portion of the HR-rich phase is 2 μm or less, particularly 1 μm or less. It is preferable. When the thickness of the thinnest part of the HR rich phase is less than 0.01 μm, the effect of increasing the coercive force may not be sufficient, and when the thickness of the thickest part of the HR rich phase exceeds 2 μm, the residual magnetic flux density Br may decrease.

In the HR-rich phase, HR replaces the occupied site of R. The content of Nd in the HR-rich phase is 0.8 times (80%) or less, particularly 0.75 times (75%) or less, and particularly 0.7 times (70%) the content of Nd in the central portion of the main phase. %) or less is preferable. If this ratio exceeds the above range, the effect of increasing the coercive force by HR may not be sufficient.

In the HR-rich phase, the area of the HR-rich phase evaluated in a cross section within 200 μm from the surface of the sintered magnet (for example, the diffusion surface in the grain boundary diffusion treatment described later) is 2 with respect to the area of the entire main phase. % Or more, particularly 4% or more, and particularly preferably 5% or more. If the ratio of the HR rich phase is less than the above range, the effect of increasing the coercive force may not be sufficient. The proportion of this HR-rich phase is preferably 40% or less, particularly 30% or less, and especially 25% or less. If the ratio of the HR rich phase exceeds the above range, the residual magnetic flux density Br may decrease.

Furthermore, the content of HR in the HR-rich phase is 1.5 times (150%) or more, particularly 2 times (200%) or more, and especially 3 times (300%) the content of HR in the central portion of the main phase. The above is preferable. If this ratio is below the above range, the effect of increasing the coercive force may not be sufficient.

On the other hand, the content of HR with respect to the sum of R'and HR in the HR-rich phase is preferably 20 atom% or more, particularly 25 atom% or more, and particularly 30 atom% or more. This range is more preferably more than 30 atomic %, particularly 31 atomic% or more. If the content of HR in the HR rich phase is less than the above range, the effect of increasing the coercive force may not be sufficient.

Further, the structure of the magnet of the present invention includes a grain boundary phase formed between crystal grains of the main phase. The grain boundary phase contains the (R′,HR)-Fe(Co)-M ₁ phase. Here, the case not containing Co can be expressed as (R′,HR)-Fe-M ₁ , and the case containing Co can be expressed as (R′,HR)-FeCo-M ₁ .

The grain boundary phase, (R ', HR) -M ₁ phase, preferably R' and HR total 50 atomic% or more of (R ', HR) -M ₁ phase and, like M ₂ boride phase It may be contained, and it is particularly preferable that the M ₂ boride phase be present at the grain boundary triple point. Furthermore, in the structure of the magnet of the present invention, the grain boundary phase may include an R-rich phase or an (R′,HR)-rich phase, and an R carbide or an (R′,HR) carbide or an R oxide. Compound or (R′,HR) oxide, R nitride or (R′,HR) nitride, R halide or (R′,HR) halide, R acid halide or (R′,HR) acid The phase of the compound of unavoidable impurities mixed in during the manufacturing process such as a halide may be contained, but at least the grain boundary triple point, preferably the whole intergranular grain boundary and the grain boundary triple point (the whole grain boundary phase). ), R ₂ (Fe,(Co)) ₁₇ phase or (R′,HR) ₂ (Fe,(Co)) ₁₇ phase, R _1.1 (Fe,(Co)) ₄ B ₄ phase or (R′, It is preferred that the HR) _1.1 (Fe, (Co)) ₄ B ₄ phase not be present.

The grain boundary phase is formed outside the crystal grains of the main phase, and the (R′,HR)-Fe(Co)-M ₁ phase is preferably present in the structure of the magnet in a volume ratio of 1% or more. If the (R′,HR)-Fe(Co)-M ₁ phase is less than 1% by volume, a sufficiently high coercive force may not be obtained. The volume ratio of the (R′,HR)-Fe(Co)-M ₁ phase is preferably 20% or less, particularly preferably 10% or less. When the (R′,HR)-Fe(Co)-M ₁ phase exceeds 20% by volume, the residual magnetic flux density Br may be significantly reduced.

The (R′,HR)-Fe(Co)-M ₁ phase is a phase of Fe alone when it does not contain Co, and Fe and a compound containing Co when it contains Co, and has a space group I4/mcm. It is considered to be a phase of an intermetallic compound having a crystal structure of, for example, (R′,HR) ₆ (Fe,(Co)) ₁₃ Si phase, (R′,HR) ₆ (Fe,(Co)) ₁₃ Ga phase, (R′,HR) ₆ (Fe,(Co)) ₁₃ Al phase, and other (R′,HR) ₆ (Fe,(Co)) ₁₃ (M ₁ ) phases. The (R′,HR)-Fe(Co)-M ₁ phase is distributed so as to surround the crystal grains of the main phase, so that the adjacent main phases are magnetically separated, and as a result, the coercive force is improved.

It can be said that the (R′,HR)-Fe(Co)-M ₁ phase is a phase in which a part of R is HR in the phase represented by R-Fe(Co)-M ₁ . (R content of HR for and total HR ', HR) -Fe (Co ) R ₁ phase -M' is preferably not more than 30 atomic%. In general, R-Fe (Co) -M 1 phase, La, Pr, it is possible to form the compound phase, such light rare earth and a stable as Nd, such as a portion of the rare earth element is Dy, Tb and Ho When it is replaced with a heavy rare earth element (HR), a stable phase is formed up to the above HR content of 30 atomic %. If it exceeds this, a ferromagnetic phase such as the (R′,HR) ₁ Fe ₃ phase is generated in the low temperature heat treatment step described later, for example, and the coercive force and the squareness may be deteriorated. The lower limit of this range is not particularly limited, but is usually 0.1 atom% or more.

Incidentally, M ₁ in the (R ', HR) -Fe ( Co) -M 1 phase,
0.5 to 50 atomic% is Si, and the balance is selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. Be one or more elements,
1.0 to 80 atomic% is Ga, and the rest is Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements, or 0.5 to 50 atomic% of Al and the balance of Si, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, It is preferably one or more elements selected from Au, Hg, Pb and Bi.

These elements are the above-mentioned intermetallic compound ((R',HR) ₆ (Fe,(Co)) ₁₃ Si phase, (R',HR) ₆ (Fe,(Co)) ₁₃ Ga phase, (R', HR) ₆ (Fe, (Co)) ₁₃ Al phase and the like (R′, HR) ₆ (Fe, (Co)) ₁₃ (M ₁ ) phase and the like are stably formed, and the M ₁ sites are mutually formed. Can be replaced with Even if the elements of the M ₁ site are compounded, no significant difference is found in the magnetic characteristics, but in practice, the variation in the magnetic characteristics is reduced to stabilize the quality, and the cost is reduced by reducing the amount of expensive elements added. Planned.

In the R-Fe-B system sintered magnet of the present invention, it is preferable that the grain boundary phase is distributed so as to individually surround the crystal grains of the main phase at the inter-grain boundary and the grain boundary triple point. , That each grain is separated from other neighboring grains by the grain boundary phase, for example, when focusing on each grain, if the grain is the core, the grain boundary phase forms a crystal as a shell. It is more preferable to have a structure covering the grains (structure similar to so-called core/shell structure). As a result, the crystal grains of the adjacent main phase are magnetically separated, and the coercive force is further improved. In order to ensure magnetic separation of the main phase, the thickness of the narrowest part of the grain boundary phase sandwiched between two adjacent crystal grains is preferably 10 nm or more, particularly preferably 20 nm or more, and 500 nm or less. More preferably, it is particularly preferably 300 nm or less. When the width of the grain boundary phase is narrower than 10 nm, there is a possibility that a sufficient coercive force improving effect due to magnetic separation cannot be obtained. Further, the average thickness of the narrowest part of the grain boundary phase sandwiched between two adjacent crystal grains is preferably 50 nm or more, particularly 60 nm or more, and more preferably 300 nm or less, particularly 200 nm or less.

The coverage of the surface of the crystal grains of the main phase with the grain boundary phase is preferably 50% or more, particularly 60% or more, and especially 70% or more, and the entire surface of the crystal grains may be coated. The rest of the grain boundary phase may be, for example, a (R′,HR)-M ₁ phase in which the sum of R′ and HR is 50 atomic% or more, a M ₂ boride phase, or the like.

The grain boundary phase contains 25 to 35 atomic% R, 2 to 8 atomic% M ₁ , 8 atomic% or less (that is, 0 atomic% or more than 0 atomic% and 8 atomic% or less) Co, and the balance. It is preferable that the (R′,HR)-Fe(Co)-M ₁ phase having the composition of Fe is included. This composition can be quantified by an electron probe microanalyzer (EPMA) or the like. The M ₁ sites can be mutually substituted by plural kinds of elements. The (R′,HR)-Fe(Co)-M ₁ phase is preferably present in the form of one or both of an amorphous phase and a microcrystalline phase having a grain size of 10 nm or less, preferably less than 10 nm. When the crystallization of the (R′,HR)-Fe(Co)-M ₁ phase proceeds, the (R′,HR)-Fe(Co)-M ₁ phase aggregates at the grain boundary triple points, and as a result, There is a risk that the coercive force of the magnet will decrease due to the narrow width of the intergranular phase and the discontinuity. In addition, as the crystallization of the (R′,HR)-Fe(Co)-M ₁ phase progresses, the R-rich phase or the (R′,HR)-rich phase becomes a main phase crystal grain and a grain boundary phase. Although it may be generated at the interface, the coercive force is not significantly improved by the formation of the R-rich phase or the (R′,HR)-rich phase itself.

On the other hand, when the (R′,HR)-M ₁ phase and the M ₂ boride phase are present, these phases also include one or both of the amorphous phase and the microcrystalline phase having a grain size of 10 nm or less, preferably less than 10 nm. It is preferably present in the form.

Next, a method for producing the R-Fe-B system sintered magnet of the present invention will be described below.
Each step in the production of the R-Fe-B system sintered magnet is basically the same as the ordinary powder metallurgy method, and a step of preparing an alloy fine powder having a predetermined composition (the raw material is used in this step). A melting step of melting to obtain a raw material alloy and a pulverizing step of pulverizing the raw material alloy), a step of compacting and molding an alloy fine powder in a magnetic field to obtain a compact, a sintered compact of a compact And a cooling step after sintering.

In the melting step, a predetermined composition, for example, 12 to 17 atomic% of R (R is one or more elements selected from rare earth elements including Y, and Nd is essential, and preferably Pr is contained. ), 0.1 to 3 atomic% of M ₁ (M ₁ is Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, One or more elements selected from Pb and Bi), 0.05 to 0.5 atomic% of M ₂ (M ₂ is 1 selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W) Or more elements), (4.8+2×m to 5.9+2×m) atomic% (m is the content (atomic %) of the element represented by M ₂ ) B, 10 atomic% or less Co, 0 0.5% by atom or less of C, 1.5% by atom or less of O, 0.5% by atom or less of N, and the balance of Fe, usually a raw material according to a composition not containing C, O and N. The metal or alloy of (1) is weighed, and the raw material is manufactured by melting the raw material by, for example, high-frequency melting in a vacuum or in an inert gas atmosphere, preferably an inert gas atmosphere such as Ar gas, and cooling. In the composition of the metal or alloy of this raw material, R may or may not contain at least one element (HR) selected from Dy, Tb and Ho. The raw material alloy may be cast by a normal melting casting method in which it is cast in a flat mold or a book mold, or by a strip casting method. When the α-Fe primary crystal remains in the cast alloy, this alloy is heat-treated at 700 to 1,200° C. for 1 hour or more in vacuum or in an inert gas atmosphere such as Ar gas to form a fine structure. It is possible to homogenize and eliminate the α-Fe phase.

The crushing step is a coarse crushing step in which a raw material alloy is mechanically crushed by a brown mill or the like, hydrogenated, etc., and preferably once crushed to an average particle size of 0.05 mm or more and 3 mm or less, particularly 1.5 mm or less. To do. In the case of an alloy produced by strip casting, this coarse pulverization step is preferably hydrogenation pulverization. After the coarse pulverization, further, by a fine pulverization process such as jet mill pulverization using high pressure nitrogen, etc., preferably the average particle size is 0.2 μm or more, particularly 0.5 μm or more, 30 μm or less, particularly 20 μm or less, and especially An alloy fine powder having a size of 10 μm or less is manufactured. Incidentally, an additive such as a lubricant may be added, if necessary, in one or both of the steps of coarse pulverization and fine pulverization of the raw material alloy.

The two-alloy method may be applied to the production of the alloy fine powder. According to this method, a mother alloy having a composition close to R ₂ —T ₁₄ —B ₁ (T usually represents Fe or Fe and Co) and a sintering aid alloy having a rare earth (R)-rich composition are respectively prepared. It is manufactured, coarsely crushed, and then the mixed powder of the mother alloy and the sintering aid obtained is crushed by the above method. The above-mentioned casting method or melt-span method may be employed to obtain the sintering aid alloy.

In the forming step, while finely pulverized alloy fine powder is being applied with a magnetic field, for example, while applying a magnetic field of 5 kOe (398 kA/m) to 20 kOe (1,592 kA/m), the easy axis direction of magnetization of the alloy powder is oriented. , Compacting with a compression molding machine. The forming is preferably performed in a vacuum, a nitrogen gas atmosphere, an inert gas atmosphere such as Ar gas, or the like in order to suppress the oxidation of the fine alloy powder. In the sintering step, the molded body obtained in the molding step is sintered. The sintering temperature is 900° C. or higher, particularly 1,000° C. or higher, especially 1,050° C. or higher, and 1,250° C. or lower, particularly 1,150° C. or lower, particularly 1,100° C. or lower, and the sintering time is , Usually 0.5 to 5 hours. The sintered body after sintering is cooled to a temperature of preferably 400° C. or lower, more preferably 300° C. or lower, and further preferably 200° C. or lower. The cooling rate is not particularly limited, but is preferably 1° C./minute or higher, particularly preferably 5° C./minute or higher, more preferably 100° C./minute or lower, and particularly preferably 50° C./minute or lower until the upper limit of the above range is reached. .. The sintered body after sintering is subjected to an aging treatment (for example, at 400 to 600° C. for 0.5 to 50 hours) as necessary, and then usually cooled to room temperature.

Here, a heat treatment step may be performed on the obtained sintered body (sintered magnet body). In this heat treatment step, the sintered body cooled to a temperature of 400° C. or lower is heated to 700° C. or higher, particularly 800° C. or higher, 1,100° C. or lower, particularly 1,050° C. or lower, and is heated to 400° C. or lower. After the high temperature heat treatment step of cooling again and the high temperature heat treatment step, a two-stage heat treatment step of heating at a temperature in the range of 400 to 600° C. and cooling to 300° C. or lower, preferably 200° C. or lower is performed. Preferably. The heat treatment atmosphere is preferably vacuum or an inert gas atmosphere such as Ar gas.

The heating rate of the high temperature heat treatment is not particularly limited, but is preferably 1° C./min or more, particularly 2° C./min or more, and 20° C./min or less, particularly 10° C./min or less. The holding time at the high temperature heat treatment temperature is preferably 1 hour or longer, usually 10 hours or shorter, and preferably 5 hours or shorter. After heating, it is cooled to preferably 400°C or lower, more preferably 300°C or lower, and further preferably 200°C or lower. The cooling rate is not particularly limited, but is preferably 1° C./minute or higher, particularly preferably 5° C./minute or higher, more preferably 100° C./minute or lower, and particularly preferably 50° C./minute or lower until the upper limit of the above range is reached. ..

In the low temperature heat treatment subsequent to the high temperature heat treatment, the cooled sintered body is heated at a temperature in the range of 400°C or higher, preferably 450°C or higher and 600°C or lower, preferably 550°C or lower. The heating rate of the low temperature heat treatment is not particularly limited, but is preferably 1° C./min or more, particularly 2° C./min or more, and 20° C./min or less, particularly 10° C./min or less. The holding time after heating at the low temperature heat treatment temperature is 0.5 hours or more, particularly 1 hour or more, and preferably 50 hours or less, particularly preferably 20 hours or less. The cooling rate after heating is not particularly limited, but is preferably 1° C./min or more, particularly preferably 5° C./min or more, 100° C./min or less, particularly 80° C./min or less until reaching the upper limit of the above range. Particularly, 50° C./min or less is more preferable. After the heat treatment, the sintered body is usually cooled to room temperature.

Various conditions in the high-temperature heat treatment and the low-temperature heat treatment include the composition such as the type and content of the M ₁ element, the concentration of impurities, particularly the concentration of impurities caused by the atmospheric gas at the time of manufacturing, the sintering conditions, and the like. It can be appropriately adjusted within the above-mentioned range according to the variation caused by the manufacturing process other than the heat treatment.

In the present invention, the HR rich phase containing the (R′,HR) ₂ (Fe,(Co)) ₁₄ B phase and the grain boundary phase containing the (R′,HR)—Fe(Co)—M ₁ phase are It can be formed by the field diffusion method. As the grain boundary diffusion treatment, the sintered body is processed into, for example, a sintered body having a predetermined shape and size close to the final product shape by cutting or surface grinding, and then, for example, HR. The surface of the sintered body is a metal, compound or intermetallic compound containing an element (HR element) represented by (HR is one or more elements selected from Dy, Tb and Ho), for example, in the form of powder or thin film. Can be applied to the inside of the sintered body from the surface of the sintered body through the grain boundary phase from the metal, compound or intermetallic compound that surrounds the sintered body. It should be noted that an element represented by HR may be solid-solved in a portion other than the HR-rich phase of the main phase due to grain boundary diffusion, but in particular, it should not be solid-soluted in the central portion of the main phase. Is preferred. On the other hand, it is preferable that no rare earth element other than the element represented by HR forms a solid solution in the main phase due to grain boundary diffusion.

As a grain boundary diffusion method of introducing the HR element from the surface of the sintered body into the sintered body through the grain boundary phase,
(1) A method in which a powder containing a metal, a compound, or an intermetallic compound containing an HR element is placed on the surface of a sintered body and heat-treated in a vacuum or an inert gas atmosphere (for example, a dip coating method),
(2) A method of forming a thin film of a metal, a compound, or an intermetallic compound containing an HR element on the surface of a sintered body in a high vacuum, and performing heat treatment in a vacuum or in an inert gas atmosphere (for example, a sputtering method),
(3) A metal, compound or intermetallic compound containing an HR element is heated in a high vacuum to form a vapor phase containing the HR element, and the HR element is supplied and diffused into the sintered body through the vapor phase. Method (eg vapor diffusion method)
And the like, and the method of (1) or (2), especially (1) is particularly preferable.

Suitable metal, compound or intermetallic compound containing HR element includes, for example, elemental metal or alloy of HR element, oxide of HR element, halide, acid halide, hydroxide, carbide, carbonate, nitride. Compounds, hydrides, borides, and mixtures thereof, intermetallic compounds of HR elements and transition metals such as Fe, Co, and Ni (a part of the transition metals is Si, Al, Ti, V, Cr, Mn, Cu. , Zn, Ga, Ge, Pd, Ag, Cd, Zr, Nb, Mo, In, Sn, Sb, Hf, Ta, W, Pt, Au, Hg, Pb and Bi. Can also be).

The thickness of the HR rich phase can be controlled by adjusting the addition amount of the HR element, the diffusion amount of the HR element into the sintered body, or the treatment temperature or treatment time in the grain boundary diffusion treatment.

In the present invention, the HR rich phase containing the (R′,HR) ₂ (Fe,(Co)) ₁₄ B phase and the grain boundary phase containing the (R′,HR)—Fe(Co)—M ₁ phase are used as the grain boundaries. In order to form by diffusion, on the surface of the sintered body cooled after the sintering or after the heat treatment before the grain boundary diffusion treatment, a metal, a compound or an intermetallic compound, for example, is arranged in the form of a powder or a thin film, first, Grain boundary diffusion of the HR element into the sintered body by heating to a temperature higher than 950°C, particularly 960°C or higher, especially 975°C or higher, and 1,100°C or lower, particularly 1,050°C or lower, especially 1,030°C or lower. Then, a high temperature heat treatment step of cooling to 400° C. or lower, preferably 300° C. or lower, and more preferably 200° C. or lower is performed. The heat treatment atmosphere is preferably vacuum or an inert gas atmosphere such as Ar gas.

If the heating temperature is lower than the above range, the effect of increasing the coercive force may not be sufficient, and if it exceeds the above range, the coercive force may decrease due to grain growth. Further, it is effective that the heating temperature is equal to or higher than the peritectic temperature (decomposition temperature) of the (R′,HR)-Fe(Co)-M ₁ phase. (R ', HR) -Fe ( Co) -M 1 phase, depending on the type of M _1, high temperature stability is changed, depending on the type of M _1, (R', HR) -Fe ( Co) -M 1 The peritectic temperatures forming the phases are different. The peritectic temperature is, for example, 640° C. when M ₁ is Cu, 750° C. when M ₁ is Al, 850° C. when M ₁ is Ga, 890° C. when M ₁ is Si, and M ₁ is The temperature is 960° C. for Ge and 890° C. for M ₁ is In. The heating rate in this case is not particularly limited, but is preferably 1° C./min or more, particularly 2° C./min or more, 20° C./min or less, 10° C./min or less. The heating time is 0.5 hours or longer, particularly 1 hour or longer, preferably 50 hours or shorter, particularly preferably 20 hours or shorter.

The cooling rate after heating is not particularly limited, but is preferably 1° C./min or more, particularly 5° C./min or more, 100° C./min or less, and particularly 50° C./min or less until the upper limit of the above range is reached. More preferable. If the cooling rate is less than the above range, (R ', HR) for -Fe (Co) -M ₁ phase is segregated at the grain boundary triple point, there is a possibility that magnetic characteristics deteriorate. On the other hand, when the cooling rate exceeds 100° C./min, segregation of the (R′,HR)-Fe(Co)-M ₁ phase in the cooling process can be suppressed, but the squareness of the sintered magnet is particularly high. May get worse.

After the high temperature heat treatment, a low temperature heat treatment step of heating at 400° C. or higher, particularly 430° C. or higher, to 600° C. or lower, particularly 550° C. or lower, and then cooling to a temperature of 300° C. or lower, preferably 200° C. or lower is performed. carry out. The heat treatment atmosphere is preferably vacuum or an inert gas atmosphere such as Ar gas.

The heating temperature is, the grain boundary phase (R ', HR) -Fe ( Co) to form a -M ₁ phase, (R' and, HR) -Fe (Co) ₁ phase of less than the peritectic temperature -M It is effective to do. If this heating temperature is lower than 400° C., the reaction rate of forming (R′,HR)—Fe(Co)—M ₁ may be very slow. On the other hand, when the heating temperature exceeds 600° C., the reaction rate of forming (R′,HR)-Fe(Co)-M ₁ is very fast, and (R′,HR)-Fe(Co)-M ₁ particles The boundary phase may be largely segregated at the grain boundary triple points, and the magnetic properties may be deteriorated.

The heating rate of the low temperature heat treatment is not particularly limited, but is preferably 1° C./min or more, particularly 2° C./min or more, and 20° C./min or less, particularly 10° C./min or less. The holding time after heating at the low temperature heat treatment temperature is 0.5 hours or more, particularly 1 hour or more, and preferably 50 hours or less, particularly preferably 20 hours or less. The cooling rate after heating is not particularly limited, but is preferably 1° C./min or more, particularly preferably 5° C./min or more, 100° C./min or less, particularly 80° C./min or less until reaching the upper limit of the above range. Particularly, 50° C./min or less is more preferable. After the heat treatment, the sintered body is usually cooled to room temperature.

Hereinafter, the present invention will be specifically described with reference to Reference Examples, Examples and Comparative Examples, but the present invention is not limited to the following Examples.

[Reference Examples 1 and 2]
Rare earth metal (Nd or didymium (mixture of Nd and Pr)), electrolytic iron, Co, metal or alloy of M ₁ element and M ₂ element, and Fe-B alloy (ferroboron) to obtain a predetermined composition. Thus, the alloy ribbon was manufactured by melting in an Ar gas atmosphere in a high frequency induction furnace and strip casting the molten alloy on a water-cooled copper roll. The thickness of the obtained alloy ribbon was about 0.2 to 0.3 mm.

Next, the produced alloy ribbon was subjected to hydrogen occlusion treatment at room temperature and then heated at 600° C. in vacuum for dehydrogenation to powder the alloy. 0.07% by mass of stearic acid as a lubricant was added to and mixed with the obtained coarse powder. Next, the mixture of the coarse powder and the lubricant was pulverized by a jet mill in a nitrogen stream to produce fine powder having an average particle size of about 3 μm.

Next, the prepared fine powder was filled in a mold of a molding apparatus in an inert gas atmosphere, and while being oriented in a magnetic field of 15 kOe (1.19 MA/m), pressure molding was performed in a direction perpendicular to the magnetic field. did. Next, the obtained green compact is sintered in vacuum at 1,050 to 1,100° C. for 3 hours, cooled to 200° C. or lower, and then aged at 450 to 530° C. for 2 hours. A sintered body (sintered magnet body) was obtained. The composition of the obtained sintered body is shown in Table 1, and its magnetic properties are shown in Table 2. The magnetic characteristics were evaluated by cutting the center of the obtained sintered body into a rectangular parallelepiped shape having a size of 6 mm×6 mm×2 mm.

[Examples 1 to 6, Comparative Examples 1 to 3]
The sintered body obtained in Reference Example 1 was processed into a rectangular parallelepiped shape having a size of 20 mm×20 mm×2.2 mm, and then terbium oxide particles having an average particle diameter of 0.5 μm were mixed with ethanol in a mass fraction of 50%. It was dipped and dried to form a coating film of terbium oxide on the surface of the sintered body. Next, the sintered body on which the coating film was formed was heated in vacuum at the holding temperature and holding time shown in Table 2 and then subjected to high temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2. Then, after heating for 2 hours at the holding temperature shown in Table 2, low temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2 was performed to obtain a sintered magnet. The composition of the obtained sintered magnet is shown in Table 1, and its magnetic properties are shown in Table 2. The magnetic characteristics were evaluated by cutting out the center of the obtained sintered magnet into a rectangular parallelepiped shape having a size of 6 mm×6 mm×2 mm.

FIG. 1 is an element distribution of Nd (A) and an element distribution of Tb (B) obtained by observing the element distribution within 200 μm from the diffusion surface of the sintered magnet manufactured in Example 2 with an electron probe microanalyzer (EPMA). ) Image. It can be seen that Tb is diffused through the grain boundary phase and the HR rich phase is nonuniformly generated on the surface of the main phase. This HR-rich phase was confirmed to be the (R′,HR) ₂ (Fe,(Co)) ₁₄ B phase, and it existed at the inter-grain grain boundary and the grain boundary triple point, and in particular, the grain boundary triple point. Was thickly present in the part. On the other hand, it was confirmed that the grain boundary phase contained the (R′,HR)-Fe(Co)-M ₁ phase and the (R′,HR) rich phase, and further (R′,HR). It was confirmed that the oxide phase was segregated mainly at the grain boundary triple points.

On the other hand, FIG. 2 shows the element distribution within 200 μm from the diffusion surface of the sintered magnet prepared in Comparative Example 2 observed with an electron probe microanalyzer (EPMA), the element distribution of Nd (A) and the element distribution of Tb. It is an image of (B). Also in this case, it can be seen that Tb diffuses through the grain boundary phase and the HR rich phase is generated on the surface portion of the main phase, but the HR rich phase is uniformly generated on the surface portion of the main phase. ing.

In the image of the Tb element distribution, the Tb element distribution does not clearly distinguish the HR-rich phase from the (R′,HR)-rich phase and the (R′,HR)oxide phase, but the Nd-element distribution. Focusing on the image of, the Nd content is higher in the (R′,HR) rich phase and the (R′,HR) oxide phase and lower in the HR rich phase than in the central portion of the main phase, It is possible to distinguish. In the cross sections of the R-Fe-B system sintered magnets of Examples and Comparative Examples, a portion having an Nd content of 80% or less compared to the Nd content in the central portion of the main phase was defined as an HR rich phase, and its partial area Table 2 shows the result of evaluation of the ratio of the main phase to the total area of the main phase. It was found that the ratio of the HR rich phase of the sintered magnet of the example was higher than that of the sintered magnet of the comparative example, and this R—Fe—B based sintered magnet had a high coercive force.

[Examples 7 to 9 and Comparative Example 4]
The sintered body obtained in Reference Example 2 was processed into a rectangular parallelepiped shape having a size of 20 mm×20 mm×2.2 mm, and then terbium oxide particles having an average particle diameter of 0.5 μm were mixed with ethanol in a mass fraction of 50%. It was dipped and dried to form a coating film of terbium oxide on the surface of the sintered body. Next, the sintered body on which the coating film was formed was heated in vacuum at the holding temperature and holding time shown in Table 2 and then subjected to high temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2. Then, after heating for 2 hours at the holding temperature shown in Table 2, low temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2 was performed to obtain a sintered magnet. The composition of the obtained sintered magnet is shown in Table 1, and its magnetic properties are shown in Table 2. The magnetic characteristics were evaluated by cutting out the center of the obtained sintered magnet into a rectangular parallelepiped shape having a size of 6 mm×6 mm×2 mm. Table 2 shows the ratio of the HR rich phase evaluated by the above method. It was found that the ratio of the HR rich phase of the sintered magnet of the example was higher than that of the sintered magnet of the comparative example, and this R—Fe—B based sintered magnet had a high coercive force.

[Example 10, Comparative Example 5]
The sintered body obtained in Reference Example 1 was processed into a rectangular parallelepiped shape having a size of 20 mm×20 mm×2.2 mm, and then dysprosium oxide particles having an average particle diameter of 0.5 μm were mixed with ethanol at a mass fraction of 50%. It was dipped and dried to form a coating film of dysprosium oxide on the surface of the sintered body. Next, the sintered body on which the coating film was formed was heated in vacuum at the holding temperature and holding time shown in Table 2 and then subjected to high temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2. Then, after heating for 2 hours at the holding temperature shown in Table 2, low temperature heat treatment of cooling to 200° C. at the cooling rate shown in Table 2 was performed to obtain a sintered magnet. The composition of the obtained sintered magnet is shown in Table 1, and its magnetic properties are shown in Table 2. The magnetic characteristics were evaluated by cutting out the center of the obtained sintered magnet into a rectangular parallelepiped shape having a size of 6 mm×6 mm×2 mm. Table 2 shows the ratio of the HR rich phase evaluated by the above method. It was found that the ratio of the HR rich phase of the sintered magnet of the example was higher than that of the sintered magnet of the comparative example, and this R—Fe—B based sintered magnet had a high coercive force.

Claims (7)

12 to 17 atomic% of R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic% of M ₁ (M ₁ is One or more elements selected from Si, Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi), 0.05 ˜0.5 atomic% of M ₂ (M ₂ is one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W), (4.8+2×m to 5.9+2) Xm) B of atomic% (m is the content (atomic %) of the element represented by M ₂ ), 10 atomic% or less of Co, 0.5 atomic% or less of C, and 1.5 atomic% or less of O , R of 0.5 atomic% or less, and the balance of Fe, and an R-Fe-B based sintered magnet having an R ₂ (Fe, (Co)) ₁₄ B intermetallic compound as a main phase. hand,
Formed between the main phase and the crystal grains of the main phase, 25 to 35 atomic% of (R′, HR), 2 to 8 atomic% of M ₁ , 8 atomic% or less of Co, and the balance of Fe (R′,HR)-Fe(Co)-M ₁ phase having a composition (R′ includes Y, and is one or more elements selected from rare earth elements other than Dy, Tb and Ho, and Nd. HR has a grain boundary phase containing one or both of an amorphous phase and a microcrystalline phase having a grain size of 10 nm or less, and HR has one or more elements selected from Dy, Tb and Ho.
(R ', HR) M ₁ in -Fe (Co) -M ₁ phase,
0.5 to 50 atomic% is Si, and the balance is selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements,
1.0 to 80 atomic% is Ga, and the rest is Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements, or 0.5 to 50 atomic% of Al and the balance of Si, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg. , One or more elements selected from Pb and Bi,
The main phase includes an HR rich phase represented by (R′,HR) ₂ (Fe,(Co)) ₁₄ B on the surface thereof, and the content of HR in the HR rich phase is An R-Fe-B based sintered magnet having a higher HR content in the center of the phase. The R-Fe-B system sintered magnet according to claim 1, wherein the HR-rich phase is nonuniformly formed on a surface portion of the main phase. Content of Nd in the said HR rich phase is 0.8 times or less of content of Nd in the center part of the said main phase, The R-Fe-B type|system|group baking of Claim 1 or 2 characterized by the above-mentioned. Binding magnet. The area of the HR rich phase evaluated in a cross section within 200 μm from the surface of the R-Fe-B system sintered magnet is 2% or more with respect to the area of the entire main phase. The R-Fe-B system sintered magnet according to any one of 3 above. The (R′,HR)-Fe(Co)-M ₁ phase is present in the structure of the magnet in an amount of 1% or more and 20% or less in volume ratio, according to any one of claims 1 to 4. R-Fe-B system sintered magnet. The (R ', HR) M ₁ in -Fe (Co) -M ₁ phase,
0.5 to 50 atomic% is Si, and the balance is selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi. One or more elements, or 1.0 to 80 atomic% Ga, the balance Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg The R-Fe-B system sintered magnet according to any one of claims 1 to 5, wherein the R-Fe-B based sintered magnet is one or more elements selected from the group consisting of Pb, Pb and Bi. The residual magnetic flux density Br is 13.2 kG or more at 23° C., and the coercive force Hcj is 26.2 kOe or more at 23° C. 7. The R- according to claim 1. Fe-B system sintered magnet.