JP7179799B2 - R-Fe-B system sintered magnet - Google Patents

Description

The present invention relates to an R—Fe—B system sintered magnet having high coercive force.

Nd--Fe--B based sintered magnets (hereinafter sometimes referred to as Nd magnets) are functional materials essential for energy saving and high performance, and their application range and production volume are expanding year by year. For example, in automotive applications, it is expected that the Nd magnets will be used in a high-temperature environment. At the same time, high coercive force is required. On the other hand, the coercive force of Nd magnets tends to decrease significantly 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 technique for increasing the coercive force of Nd magnets, it is effective to replace part of the Nd in the Nd ₂ Fe ₁₄ B compound, which is the main phase, with Dy or Tb. However, there is a risk that the price will be unstable and fluctuate greatly because the production areas where it is commercially viable are limited and the stable supply is affected by geopolitical factors. Against this background, in order to capture a large market for R--Fe--B magnets that can be used at high temperatures, new methods or R - Development of Fe-B magnet compositions is required. From this point of view, various methods have been conventionally proposed.

For example, in Japanese Patent No. 3997413 (Patent Document 1), the atomic percentage of R is 12 to 17% (R is at least two 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 (However, Fe is Al, Ti, V, Cr, Mn, Ni in a substitution amount of 3 atomic % or less , Cu, Zn, Ga, Ge, Zr, Nb, Mo, In, Sn, Sb, Hf, Ta, W, Pt, Au, Hg, Pb, and Bi. R—Fe—B system sintered magnet having a composition of R ₂ (Fe, (Co), Si) ₁₄ B intermetallic compound as a main phase and having a coercive force of at least 10 kOe or more, wherein B At least an R—Fe(Co)—Si grain boundary phase containing no rich phase and consisting of atomic percentages of 25 to 35% R, 2 to 8% Si, 8% or less Co, and the balance Fe An R—Fe—B system sintered magnet having 1% or more of the entire magnet is disclosed. This sintered magnet is cooled at a controlled rate of 0.1 to 5° C./min to at least 700 to 500° C. in the cooling step during sintering or during heat treatment after sintering in its manufacture. Alternatively, a R--Fe(Co)--Si grain boundary phase is formed in the structure by holding a constant temperature for at least 30 minutes during cooling and cooling in multiple stages.

Japanese National Publication of International Patent Application No. 2003-510467 (Patent Document 2) discloses an Nd--Fe--B alloy with a low boron content, a sintered magnet from this alloy, and a method for producing the same, and produces a sintered magnet from this alloy. As a method for doing so, when cooling the raw material 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) describes an RTB magnet containing an R ₂ Fe ₁₄ B main phase and a grain boundary phase. A part of this grain boundary phase is an R-rich phase containing more R than the main phase, and the other grain boundary phase is a transition metal-rich phase with a lower rare earth element concentration and a higher transition metal element concentration than the main phase. be. It also describes that this RTB rare earth sintered magnet is manufactured by performing heat treatment at 400°C to 800°C after performing sintering at 800°C to 1,200°C.

Japanese Patent Application Laid-Open No. 2014-132628 (Patent Document 4) discloses that the grain boundary phase includes an R-rich phase having a total atomic concentration of rare earth elements of 70 atomic % or more and a total atomic concentration of rare earth elements of 25 to 35 atomic %. An RTB rare earth sintered magnet containing a ferromagnetic transition metal-rich phase and having an area ratio of the transition metal-rich phase in the grain boundary phase of 40% or more is described. 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 less, and further performing the second heat treatment step at 450 ° C. C. to 600.degree. C. in a plurality of heat treatment processes.

Japanese Patent Application Laid-Open No. 2014-146788 (Patent Document 5) describes an RTB rare earth sintered magnet having a main phase composed of R ₂ Fe ₁₄ B and a grain boundary phase containing more R than the main phase. , the axis of easy magnetization of the R ₂ Fe ₁₄ B main phase is parallel to the c-axis, the crystal grain shape of the R ₂ Fe ₁₄ B main phase is an elliptical shape extending in a direction orthogonal to the c-axis direction, and the grain boundary phase However, an RTB rare earth sintered magnet containing an R-rich phase having a total atomic concentration of rare earth elements of 70 atomic % or more and a transition metal rich phase having a total atomic concentration of rare earth elements of 25 to 35 atomic % is described. It also describes that in the production thereof, 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 Patent Laying-Open No. 2014-209546 (Patent Document 6) discloses a R ₂ T ₁₄ B main phase and a two-particle grain boundary phase between crystal grains of two adjacent R ₂ T ₁₄ B main phases. A rare earth magnet is disclosed in which the two grain boundary phase has a thickness of 5 nm or more and 500 nm or less and is composed of a phase having magnetism different from that of a ferromagnetic material. This rare earth magnet is formed from a compound that does not become ferromagnetic while containing the T element as a two-grain boundary phase, so that this phase contains transition metal elements such as Al, Ge, Si, It contains M elements such as Sn and Ga. Furthermore, 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 boundary phase, and a La ₆ Co ₁₁ Ga ₃ type two-grain boundary phase can be formed. and the crystal grains of the R ₂ T ₁₄ B main phase. It is described that it can suppress the generation of nuclei. In the production thereof, it is described that 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), a Nd ₂ Fe ₁₄ B-type compound is used as a main phase, surrounded by two main phases, and having a thickness of An RTB based sintered magnet having two grain boundaries of 5 to 30 nm and a grain boundary triple point surrounded by three or more main phases is disclosed.

Japanese Patent No. 3997413 Japanese Patent Publication No. 2003-510467 Japanese Patent No. 5572673 JP 2014-132628 A JP 2014-146788 A JP 2014-209546 A WO2014/157448 WO2014/157451

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

SUMMARY OF THE INVENTION It is an object of the present invention to provide a novel R--Fe--B based sintered magnet having a high coercive force.

As a result of extensive studies to solve the above problems, the present inventors have found that 12 to 17 atomic % of R (R is one or more elements selected from rare earth elements including Y, and Nd 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 selected elements), (4.8 + 2 × m to 5.9 + 2 × m) atomic % (m is the content of the element represented by M ₂ (atomic %)) B, 10 atomic % or less R ₂ (Fe, (Co)) ₁₄ B having a composition of Co, up to 0.5 atomic % C, up to 1.5 atomic % O, up to 0.5 atomic % N, and the balance Fe; An R—Fe—B based sintered magnet having an intermetallic compound as a main phase, which is formed between the main phase and crystal grains of the main phase and contains 25 to 35 atomic % of (R′, HR), 2 to (R', HR)-Fe(Co) _{-M phase} (R' contains Y, Dy, Tb and Ho and Nd is essential, and HR is one or more elements selected from Dy, Tb and Ho), an amorphous phase and fine particles having a particle size of 10 nm or less. and a grain boundary phase included in one or both of the crystal phases, and the main phase is an HR-rich phase represented by (R', HR) ₂ (Fe, (Co)) ₁₄ B at the surface portion wherein the HR content in the HR-rich phase is higher than the HR content in the center of the main phase, and the R—Fe—B sintered magnet has a high coercive force found to be magnets.

Then, such an R—Fe—B based sintered magnet has a step of preparing an alloy fine powder having a predetermined composition, a step of compacting the alloy fine powder while applying a magnetic field to obtain a compact, and a compact of 900. A step of sintering at a temperature in the range of ~1,250°C to obtain a sintered body, a step of cooling the sintered body to a temperature of 400°C or less, HR (HR is one or more selected from Dy, Tb and Ho element) is placed on the surface of the sintered body and heated at a temperature in the range of more than 950 ° C. and 1,100 ° C. or less to transfer the HR to the grain boundaries We found that it can be produced by a method including a high temperature heat treatment step of diffusing and cooling to 400 ° C. or less, and a low temperature heat treatment step of heating at a temperature in the range of 400 to 600 ° C. and cooling to 300 ° C. or less after the high temperature heat treatment. I came up with the invention.

Accordingly, the present invention provides the following R—Fe—B based sintered magnet.
[1] 12 to 17 atomic % R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic % 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 of the element represented by M ₂ (atomic %)) B, 10 atomic % or less Co, 0.5 atomic % or less C, 1.5 atomic % An R—Fe—B system sintered magnet having a composition of O below, N of 0.5 atomic percent or less, and Fe as the balance, and having an R ₂ (Fe, (Co)) ₁₄ B intermetallic compound as a main phase. An R—Fe—B based sintered magnet with
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 formed between the main phase and the crystal grains of the main phase (R', HR)-Fe(Co)-M _phase having a composition (R' contains Y and is one or more elements selected from rare earth elements excluding Dy, Tb and Ho, and Nd is essential, and HR has a grain boundary phase containing one or more elements selected from Dy, Tb and Ho),
(R', HR)-Fe(Co)-M _{1 in the M 1 phase} is
0.5 to 50 atomic % Si, the balance 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 % of Ga and the balance selected from 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 , is one or more elements selected from Pb and Bi,
The main phase contains an HR-rich phase represented by (R′, HR) ₂ (Fe, (Co)) ₁₄ B on its surface, and the content of HR in the HR-rich phase is higher than the content of HR in the core of the phase,
R--Fe-- characterized in that the area of the HR-rich phase evaluated in a cross section within 200 μm from the surface of the R--Fe--B based sintered magnet is 2% or more with respect to the area of the entire main phase. B-based sintered magnet.
[2] The R--Fe--B system sintered magnet according to [1], wherein the HR-rich phase is unevenly formed on the surface of the main phase.
[3] The R- Fe—B system sintered magnet.
[4] The R--Fe--B-based sintered magnet according to any one of [1] to [3], wherein the total content of Dy, Tb and Ho relative to the composition of the entire magnet is 1 atomic % or less. binding magnets.
[5] Any one of [1] to [4], wherein the (R', HR)-Fe(Co)-M ₁ phase is present in the magnet structure at a volume fraction of 1% or more and 20% or less. 1. The R—Fe—B based sintered magnet according to 1.
[6] M ₁ in the (R', HR)-Fe(Co)-M ₁ phase is
0.5 to 50 atomic % Si, the balance 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 % of Ga and the balance of Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg , Pb and Bi.
[7] Any one of [1] to [6], wherein 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 R—Fe—B based sintered magnet described above.
[ 8 ] The above-mentioned M ₂ is one or more elements selected from V, Cr, Zr, Nb, Mo, Hf, Ta and W. Any one of [1] to [ 7 ] R—Fe—B system sintered magnet.
The present invention also relates to the following R—Fe—B based sintered magnet.
[1] 12 to 17 atomic % R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic % 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 of the element represented by M ₂ (atomic %)) B, 10 atomic % or less Co, 0.5 atomic % or less C, 1.5 atomic % R--Fe--B system sintering having the following composition of O, 0.5 atomic % or less of N, and the balance of Fe, and having an R ₂ (Fe, (Co)) ₁₄ B intermetallic compound as a main phase being a magnet,
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 formed between the main phase and the crystal grains of the main phase (R', HR)-Fe(Co)-M _phase having a composition (R' contains Y and is one or more elements selected from rare earth elements excluding Dy, Tb and Ho, and Nd and HR is one or more elements selected from Dy, Tb and Ho) in one or both of an amorphous phase and a microcrystalline phase with a grain size of 10 nm or less, and a grain boundary phase containing
The main phase contains an HR-rich phase represented by (R', HR) ₂ (Fe, (Co)) ₁₄ B on its surface,
An R—Fe—B system sintered magnet, wherein the HR content in the HR-rich phase is higher than the HR content in the center of the main phase.
[2] The R—Fe—B system sintered magnet according to [1], wherein the HR-rich phase is unevenly formed on the surface of the main phase.
[3] The R- according to [1] or [2], wherein the content of Nd in the HR-rich phase is 0.8 times or less the content of Nd in the center of the main phase. 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 based sintered magnet is 2% or more of the total area of the main phase [1] ] to [3].
[5] 12 to 17 atomic % R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic % 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 (M2 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 × m) atomic % (m is the content of the element represented by M ₂ (atomic %)) 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;
a step of compacting the alloy fine powder while applying a magnetic field to obtain a compact;
a step of sintering the compact at a temperature in the range of 900 to 1,250° C. to obtain a sintered compact;
cooling the sintered body to a temperature of 400° C. or less;
A metal, compound or intermetallic compound containing HR (HR is one or more elements selected from Dy, Tb and Ho) is placed on the surface of the sintered body, and A high-temperature heat treatment step of heating at a temperature in the range to diffuse grain boundaries of the sintered body and cooling to 400 ° C. or less, and after the high-temperature heat treatment, heating at a temperature in the range of 400 to 600 ° C. to 300 ° C. An R—Fe—B based sintered magnet obtained by a manufacturing method including a low-temperature heat treatment step of cooling to:

The R--Fe--B based sintered magnet of the present invention provides high coercive force even when the contents of Dy, Tb and Ho are small.

It is an image of the element distribution inside 200 μm from the diffusion surface of the sintered magnet produced in Example 2, observed with an electron probe microanalyzer (EPMA), (A) is the distribution of Nd, and (B) is the distribution of Tb. Show distribution. It is an image of the element distribution inside 200 μm from the diffusion surface of the sintered magnet produced in Comparative Example 2, observed with an electron probe microanalyzer (EPMA), (A) is the distribution of Nd, and (B) is the distribution of Tb. Show distribution.

The present invention will now be described in more detail.
First, the magnet composition of the present invention will be explained. The R--Fe--B based sintered magnet of the present invention contains 12 to 17 atomic % R, 0.1 to 3 atomic % M ₁ , 0.05 to 0.05 atomic %. 5 atomic % of M ₂ , (4.8+2×m to 5.9+2×m) atomic % of B (m is the content (atomic %) of the element represented by M ₂ ), 10 atomic % or less of Co, It has a composition of 0.5 atomic % or less C (carbon), 1.5 atomic % or less O (oxygen), 0.5 atomic % or less N (nitrogen), and the balance Fe, and contains inevitable impurities You can stay.

R is one or more elements selected from rare earth elements including Y, and Nd is essential. As the rare earth elements other than Nd, Pr, La, Ce, Gd, Dy, Tb and Ho are preferable, and Pr, Dy, Tb and Ho, especially Pr are preferable. The content of R is 12 to 17 atomic %, preferably 13 atomic % or more, and preferably 16 atomic % or less. When the content of R is less than 12 atomic %, the coercive force of the magnet is extremely lowered, and when it exceeds 17 atomic %, the residual magnetic flux density Br is lowered. Among R, the ratio of Nd, which is an essential component, is preferably 60 atomic % or more, particularly 70 atomic % or more, of the entire R. When one or more elements selected from Pr, La, Ce and Gd are included 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) is preferably Nd/(one or more elements selected from Pr, La, Ce and Gd) of 75/25 or more and 85/15 or less. In particular, when Pr is used as a rare earth element other than Nd, didymium, which is a mixture of Nd and Pr, can be used. 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 atomic % or less, more preferably 10 atomic % or less, and even more preferably 10 atomic % or less, 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 in the composition of the entire magnet is preferably 3 atomic % or less, more preferably 1.5 atomic % or less, still more preferably 1 atomic % or less, and particularly preferably It is 0.4 atomic % or less, preferably 0.01 atomic % or more. When one or more elements selected from Dy, Tb and Ho are diffused by grain boundary diffusion, the diffusion amount is preferably 0.7 atomic % or less, particularly preferably 0.4 atomic % or less, and 0.05 atomic % 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; Configured. M ₁ is _an element necessary for forming the (R', HR)-Fe(Co)-M ₁ phase described later. (Co)-M ₁ phase can be stably formed. The content of M ₁ is 0.1 to 3 atomic %, preferably 0.5 atomic % or more, and preferably 2.5 atomic % or less. If the content of M ₁ is less than 0.1 atomic %, the proportion of the (R′, HR)—Fe(Co)—M ₁ phase in the grain boundary phase is too low, and the coercive force cannot be sufficiently improved. 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 one or more elements selected from Ti, V, Cr, Zr, Nb, Mo, Hf, Ta and W. _M2 is added as an element that stably forms borides for the purpose of suppressing abnormal grain growth during sintering. The content of M ₂ is 0.05 to 0.5 atomic percent. The addition of M ₂ enables sintering at a relatively high temperature during production, leading to improved squareness and improved magnetic properties.

The content of B (boron) is (4.8+2×m) to (5.9+ ₂ ×m) atomic % (m is the content (atomic %) of the element represented by M2, the same shall apply hereinafter), It is preferably (4.9+2×m) atomic % or more, and preferably (5.7+2×m) atomic % or less. In other words, since the content of the _M2 element in the composition of the magnet of the present invention is 0.05 to 0.5 atomic percent, the content of the B content depends on the content of the _M2 element specified within the above range. Although the ranges are different, the content of B 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 content of B is an important factor. When the content of B exceeds (5.9+2×m) atomic %, the (R′, HR)—Fe(Co)—M ₁ phase, which will be described later, is not formed at the grain boundaries, and the R _1.1 Fe ₄ B ₄ compound is formed. A phase or (R',HR) _1.1 Fe ₄ B ₄ compound phase, the 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 B content is less than (4.8+2×m) atomic %, the volume fraction 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 corrosion resistance, Fe can be substituted with Co. When Co is contained, the content of Co is preferably 10 atomic % or less, particularly 5 atomic % or less. If the Co content exceeds 10 atomic %, the coercive force may be greatly reduced. The content of Co is more preferably 10 atomic % or less, particularly 5 atomic % or less, relative 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 where Co is contained and the case where Co is not contained are included.

The content of carbon, oxygen and nitrogen is preferably as low as possible, and more preferably not contained, but contamination cannot be completely avoided in the manufacturing process. As for the contents of these elements, the C (carbon) content is 0.5 atomic % or less, particularly 0.4 atomic % or less, and the O (oxygen) content is 1.5 atomic % or less, particularly 1.2 atomic % or less. atomic % or less, and the N (nitrogen) content is permissible up to 0.5 atomic % or less, particularly 0.3 atomic % or less. The content of Fe, which is the balance of the composition of the entire magnet, is preferably 70 atomic % or more, especially 75 atomic % or more, and 85 atomic % or less, especially 80 atomic % or less.

In addition to these elements, the content of elements such as H, F, Mg, P, S, Cl, Ca, etc., as inevitable impurities 0.1% by mass or less is allowed, but it is preferable that the content of these unavoidable impurities is as small as possible.

The average grain size of the R—Fe—B based 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, particularly 2 μm or more. preferable. The average grain size can be controlled by adjusting the average grain size of the fine alloy powder during 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 etchant such as Vilera's solution (for example, a mixture of glycerin:nitric acid:hydrochloric acid=3:1:2) to remove grain boundaries. A cross section of the selectively etched phase is observed with a laser microscope. Next, based on the observed image obtained, the cross-sectional area of each particle is measured by image analysis, and the diameter as an equivalent circle is calculated. Then, based on the data of the area fraction occupied by each particle size, the average diameter is obtained. Note that the average diameter may be the average of about 2,000 particles in total in 20 different images, for example.

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

The structure of the magnet of the present invention includes a main phase (crystal grains) and a grain boundary phase. The main phase that constitutes the crystal grains includes the R ₂ (Fe, (Co)) ₁₄ B intermetallic compound phase. Here, the case containing no Co can be expressed as R ₂ Fe ₁₄ B, and the 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 includes (R', HR) ₂ (Fe, (Co)) ₁₄ B (R' contains Y and is selected from rare earth elements excluding Dy, Tb and Ho). or two or more elements, Nd being essential (the same shall apply hereinafter), and HR being one or more elements selected from Dy, Tb and Ho (the same shall apply hereinafter). Here, the case containing no Co can be expressed as (R', HR) ₂ Fe ₁₄ B, and the case containing Co can be expressed as (R', HR) ₂ (Fe, Co) ₁₄ B. An HR-rich phase is a phase of an intermetallic compound in which the HR content is higher than the HR content in the core of the main phase. Pr, La, Ce and Gd are preferred as rare earth elements other than Nd in R', and Pr is particularly preferred. This HR-rich phase is formed on the surface of the main phase.

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

Further, the HR-rich phase may be formed only on a part of the surface of the main phase, for example, it may be formed so as to cover only a part 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, particularly 2% or more, and 40% or less, particularly 30% or less, of the grain size of the main phase crystal grains. In particular, it is preferably 25% or less.

The thickness of the thinnest portion of the HR-rich phase is preferably 0.01 μm or more, particularly 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. is preferred. If the thickness of the thinnest portion of the HR-rich phase is less than 0.01 μm, the effect of increasing the coercive force may not be sufficient. Br may decrease.

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

In addition, the HR-rich phase is 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). % or more, particularly 4% or more, especially 5% or more. If the proportion 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, especially 25% or less. If the proportion of the HR-rich phase exceeds the above range, the residual magnetic flux density Br may decrease.

Furthermore, the HR content in the HR-rich phase is 1.5 times (150%) or more, especially 2 times (200%) or more, especially 3 times (300%) or more that of the HR content in the center of the main phase. It is preferable that it is above. 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 atomic % or more, particularly 25 atomic % or more, and particularly preferably 30 atomic % or more. This range is more preferably greater than 30 atomic %, especially greater than or equal to 31 atomic %. If the HR content in the HR-rich phase is below the above range, the effect of increasing the coercive force may not be sufficient.

Furthermore, the structure of the magnet of the present invention includes grain boundary phases formed between crystal grains of the main phase. The grain boundary phase contains the (R', HR)-Fe(Co)-M ₁ phase. Here, the case containing no 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 includes (R', HR)-M ₁ phase, preferably (R', HR)-M ₁ phase in which the total of R' and HR is 50 atomic% or more, M ₂ boride phase, and the like. In particular, it is preferable that the _M2 boride phase exists at the grain boundary triple point. Furthermore, the structure of the magnet of the present invention may contain an R-rich phase or (R', HR) rich phase in the grain boundary phase, or may include R-carbides or (R', HR) carbides, R-oxidized or (R', HR) oxide, R nitride or (R', HR) nitride, R halide or (R', HR) halide, R acid halide or (R', HR) acid Although it may contain phases of compounds of unavoidable impurities mixed in during the manufacturing process such as halides, at least the grain boundary triple points, preferably the grain boundaries between two grains and the entire grain boundary triple points (the entire 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', The HR) _1.1 (Fe,(Co)) ₄ B ₄ phase is preferably absent.

The grain boundary phase is formed outside the crystal grains of the main phase, and the (R', HR)--Fe(Co)--M ₁ phase preferably exists in the magnet structure at 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 fraction of this (R', HR)-Fe(Co)-M ₁ phase is preferably 20% or less, particularly 10% or less. If the (R', HR)--Fe(Co)--M ₁ phase exceeds 20% by volume, the residual magnetic flux density Br may be greatly reduced.

(R', HR)-Fe(Co)-M ₁ phase is a phase of a compound containing Fe only when Co is not contained, and Fe and Co when Co is contained, space group I4/mcm For example, (R', HR) ₆ (Fe, (Co)) ₁₃ Si phase, (R', HR) ₆ (Fe, (Co)) (R', HR) ₆ (Fe, (Co)) ₁₃ (M ₁ ) phase such as ₁₃ Ga phase and (R', HR) ₆ (Fe, (Co)) ₁₃ Al phase. The (R', HR)--Fe(Co)--M ₁ phase surrounds and distributes the crystal grains of the main phase, thereby magnetically separating the adjacent main phases, thereby improving the coercive force.

The (R', HR)--Fe(Co)--M ₁ phase can be said to be a phase represented by R--Fe(Co)--M ₁ in which part of R is HR. The content of HR with respect to the sum of R' and HR in the (R', HR)--Fe(Co)--M ₁ phase is preferably 30 atomic % or less. In general, the R—Fe(Co)—M ₁ phase can form stable compound phases with light rare earth elements such as La, Pr, and Nd, and some of the rare earth elements are Dy, Tb, and Ho. Substitution with heavy rare earth elements (HR) forms a stable phase up to the above-mentioned content of HR up to 30 atomic %. If this value is exceeded, a ferromagnetic phase such as the (R', HR) ₁ Fe ₃ phase may be generated in the low-temperature heat treatment step described below, resulting in a decrease in coercive force and squareness. Although the lower limit of this range is not particularly limited, it is usually 0.1 atomic % or more.

Note that M ₁ in the (R', HR)-Fe(Co)-M ₁ phase is
0.5 to 50 atomic % Si, the balance selected from Al, Mn, Ni, Cu, Zn, Ga, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg, Pb and Bi being one or more elements;
1.0 to 80 atomic % of Ga and the balance selected from 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, At least one element selected from Au, Hg, Pb and Bi is preferable.

These elements are the above-mentioned intermetallic compounds ((R', HR) ₆ (Fe, (Co)) ₁₃ Si phase, (R', HR) ₆ (Fe, (Co)) ₁₃ Ga phase, (R', (R′, HR) ₆ (Fe, (Co)) ₁₃ (M ₁ ) phase such as HR) ₆ (Fe, (Co)) ₁₃ Al phase, etc.) is stably formed, and M ₁ sites are mutually exchanged. can be replaced with Even if the elements at the M1 site are combined, there is _no noticeable difference in magnetic properties, but in practice, it is possible to stabilize quality by reducing variations in magnetic properties and reduce costs by reducing the amount of expensive elements added. planned.

In the R—Fe—B based sintered magnet of the present invention, the grain boundary phase is preferably distributed so as to individually surround the main phase crystal grains at the grain boundary between two grains and the grain boundary triple point. , that individual crystal grains are isolated from other adjacent crystal grains by the grain boundary phase, for example, when focusing on individual crystal grains, if the crystal grain is the core, the grain boundary phase is the shell of the crystal It is more preferable to have a structure covering grains (a structure similar to a so-called core/shell structure). As a result, adjacent crystal grains of the main phase are magnetically separated, and the coercive force is further improved. In order to ensure magnetic division 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 20 nm or more, and 500 nm or less. , particularly preferably 300 nm or less. If the width of the grain boundary phase is narrower than 10 nm, there is a possibility that a sufficient effect of improving the coercive force by magnetic division cannot be obtained. 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 preferably 60 nm or more, and more preferably 300 nm or less, particularly 200 nm or less.

The grain boundary phase preferably covers the surfaces of the crystal grains of the main phase at 50% or more, particularly preferably 60% or more, particularly preferably 70% or more, and may cover the entire surface of the crystal grains. The remainder of the grain boundary phase includes, for example, (R', HR)—M ₁ phase in which the sum of R' and HR is 50 atomic % or more, M ₂ boride phase, and the like.

The grain boundary phase contains 25 to 35 atomic percent R, 2 to 8 atomic percent M ₁ , 8 atomic percent or less (i.e., 0 atomic percent or greater than 0 atomic percent to 8 atomic percent) Co, and the balance (R', HR)--Fe(Co)--M ₁ phase having an Fe composition of . This composition can be quantified using an electron probe microanalyzer (EPMA) or the like. This M ₁ site can be mutually substituted by a plurality of kinds of elements. The (R',HR)-Fe(Co)-M ₁ phase preferably exists in the form of one or both of an amorphous phase and a microcrystalline phase with a grain size of 10 nm or less, preferably less than 10 nm. As the crystallization of the (R', HR)-Fe(Co)-M _phase progresses, the (R', HR)-Fe(Co)-M _phase aggregates at the grain boundary triple points, resulting in two The width of the intergranular grain boundary phase may become narrow or discontinuous, which may reduce the coercive force of the magnet. In addition, as the crystallization of the (R', HR)--Fe(Co)--M _phase progresses, the R-rich phase or the (R', HR)-rich phase forms a boundary between the grains of the main phase and the grain boundary phase. Although it may be generated at the interface, the formation of the R-rich phase or (R', HR)-rich phase itself does not significantly improve the coercive force.

On the other hand, when the (R′, HR)—M ₁ phase and the M ₂ boride phase are present, these phases are also 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. It is preferably present in the form

Next, a method for producing the R—Fe—B based sintered magnet of the present invention will be described below.
Each step in the production of an R—Fe—B based sintered magnet is basically the same as a normal powder metallurgy method, and includes a step of preparing alloy fine powder having a predetermined composition (this step includes (including a melting process to obtain a raw material alloy by melting and a pulverizing process to pulverize the raw material alloy), a process of compacting the alloy fine powder in a magnetic field to obtain a molded body, and a process of sintering the molded body to obtain a sintered body. 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, preferably Pr ), 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); element above species), (4.8 + 2 × m to 5.9 + 2 × m) atomic % (m is the content of the element represented by M ₂ (atomic %)) B, 10 atomic % or less Co, 0 .5 atomic % or less C, 1.5 atomic % or less O, 0.5 atomic % or less N, and the balance Fe, usually C, O and N-free. The metal or alloy is weighed, for example, in a vacuum or an inert gas atmosphere, preferably an inert gas atmosphere such as Ar gas, the raw material is melted by, for example, high-frequency melting, and cooled to produce a raw material alloy. In the composition of the raw material metal or alloy, R may or may not contain one or more elements (HR) selected from Dy, Tb and Ho. Casting of the raw material alloy may be carried out using a normal melting casting method in which the alloy is cast into a flat mold or a book mold, or a strip casting method. When primary crystals of α-Fe remain in the cast alloy, the alloy is heat-treated at 700 to 1,200° C. for 1 hour or more in a vacuum or an inert gas atmosphere such as Ar gas to remove the microstructure. It can be homogenized and the α-Fe phase eliminated.

In the pulverization step, first, the raw material alloy is first subjected to coarse pulverization by mechanical pulverization such as a brown mill or hydrogenation pulverization, and is preferably pulverized to an average particle size of 0.05 mm or more and 3 mm or less, particularly 1.5 mm or less. do. This coarse grinding step is preferably hydrogrinding in the case of alloys produced by strip casting. After the coarse pulverization, it is further subjected to a fine pulverization step such as jet mill pulverization using high-pressure nitrogen, preferably having an average particle size of 0.2 μm or more, particularly 0.5 μm or more, and 30 μm or less, particularly 20 μm or less, especially An alloy fine powder of 10 µm or less is produced. Additives such as lubricants may be added in one or both of the steps of coarse pulverization and fine pulverization of the material alloy, if necessary.

A two-alloy process may be applied to produce the alloy fine powder. In this method, a master 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 composition rich in rare earth (R) are prepared. The mixed powder of the master alloy and the sintering aid is then pulverized by the above method. In order to obtain the sintering additive alloy, the casting method or the melt spun method can be employed.

In the molding step, the finely pulverized alloy fine powder is subjected to magnetic field application of, for example, 5 kOe (398 kA/m) to 20 kOe (1,592 kA/m), while orienting the easy magnetization axis direction of the alloy powder. , compacted in a compression molding machine. In order to suppress the oxidation of the fine alloy powder, the molding is preferably performed in a vacuum, a nitrogen gas atmosphere, an inert gas atmosphere such as Ar gas, or the like. In the sintering step, the compact obtained in the forming step is sintered. The sintering temperature is preferably 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, especially 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 still more preferably 200° C. or lower. The cooling rate is not particularly limited, but is preferably 1°C/min or more, particularly 5°C/min or more, and more preferably 100°C/min or less, particularly 50°C/min or less, until the upper limit of the above range is reached. . The sintered body after sintering is optionally subjected to aging treatment (for example, at 400 to 600° C. for 0.5 to 50 hours), and then usually cooled to room temperature.

Here, the obtained sintered body (sintered magnet body) may be subjected to a heat treatment step. 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, and 1,100 ° C. or lower, particularly 1,050 ° C. or lower, and heated to 400 ° C. or lower. A two-step heat treatment process is carried out: a high temperature heat treatment step of cooling again, and a low temperature heat treatment step of heating at a temperature in the range of 400 to 600 ° C. after the high temperature heat treatment step and cooling to 300 ° C. or less, preferably 200 ° C. or less. preferably. The heat treatment atmosphere is preferably a vacuum or an inert gas atmosphere such as Ar gas.

The temperature increase 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, preferably 5 hours or shorter. After heating, it is cooled to preferably 400° C. or lower, more preferably 300° C. or lower, still more preferably 200° C. or lower. The cooling rate is not particularly limited, but is preferably 1°C/min or more, particularly 5°C/min or more, and more preferably 100°C/min or less, particularly 50°C/min or less, until the upper limit of the above range is reached. .

In the low temperature heat treatment following the high temperature heat treatment, the cooled sintered body is heated to 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 rate of temperature increase in 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 raising the temperature at the low-temperature heat treatment temperature is preferably 0.5 hours or more, particularly 1 hour or more, and 50 hours or less, particularly 20 hours or less. The cooling rate after heating is not particularly limited, but is preferably 1°C/min or more, particularly 5°C/min or more, and 100°C/min or less, particularly 80°C/min or less, until the upper limit of the above range is reached. In particular, 50° C./min or less is more preferable. After the heat treatment, the sintered body is usually cooled to room temperature.

The conditions for the high-temperature heat treatment and the low-temperature heat treatment include the composition such as the _type and content of the M1 element, the concentration of impurities, especially the impurities caused by the atmosphere gas during production, and the sintering conditions. It can be appropriately adjusted within the above-described range according to fluctuations caused by manufacturing processes other than 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 a field diffusion method. As this grain boundary diffusion treatment, the sintered body is processed, for example, into a sintered body having a predetermined shape and size close to the shape of the final product by cutting, surface grinding, etc., if necessary, and then, for example, HR 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) is applied, for example, in the form of a powder or thin film to the surface of the sintered body A treatment can be applied in which the HR element is introduced from the surface of the sintered body into the interior of the sintered body through the grain boundary phase from the metal, compound, or intermetallic compound surrounding the sintered body. Note that the element represented by HR may be solid-dissolved in the main phase other than the HR-rich phase due to grain boundary diffusion, but in particular, it should not be solid-dissolved in the central part 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 be solid-dissolved in the main phase due to grain boundary diffusion.

As a grain boundary diffusion method for introducing the HR element from the surface of the sintered body into the inside of the sintered body via the grain boundary phase,
(1) A method of disposing a powder composed of a metal, a compound, or an intermetallic compound containing an HR element on the surface of a sintered body and heat-treating it in a vacuum or an inert gas atmosphere (e.g., dip coating method);
(2) A method of forming a thin film of a metal, compound or intermetallic compound containing an HR element on the surface of a sintered body in high vacuum and heat-treating it in vacuum or in an inert gas atmosphere (e.g. sputtering method);
(3) Heating a metal, compound, or intermetallic compound containing the HR element in a high vacuum to form a vapor phase containing the HR element, and supplying and diffusing the HR element into the sintered body through the vapor phase. method (e.g. vapor diffusion method)
etc., and the method (1) or (2), especially the method (1) is preferred.

Suitable metals, compounds or intermetallic compounds containing the HR element include, for example, elemental metals or alloys of the HR element, oxides, halides, acid halides, hydroxides, carbides, carbonates, nitrides of the HR element. compounds, hydrides, borides, and mixtures thereof, intermetallic compounds of HR elements and transition metals such as Fe, Co, Ni (some of the transition metals are 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 used), etc.

The thickness of the HR-rich phase can be controlled by adjusting the amount of the HR element added, the amount of the HR element diffused 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 In order to form by diffusion, a metal, a compound, or an intermetallic compound is placed, for example, in the form of a powder or a thin film on the surface of a sintered body that has been cooled after sintering or after heat treatment before grain boundary diffusion treatment, and Grain boundary diffusion of the HR element into the sintered body by heating to a temperature above 950°C, especially 960°C or higher, especially 975°C or higher, and 1,100°C or lower, especially 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, more preferably 200° C. or lower is performed. The heat treatment atmosphere is preferably a 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. Moreover, 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. The high-temperature stability of the (R', HR)-Fe(Co)-M ₁ phase varies depending on the type of M ₁ , and the (R', HR)-Fe(Co)-M ₁ phase varies depending on the type of M ₁ The peritectic temperatures that form 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 890° C. when M ₁ is Si. It is 960° C. when Ge and 890° C. when M ₁ is In. The rate of temperature increase in this case is not particularly limited, but is preferably 1° C./min or more, particularly 2° C./min or more, and 20° C./min or less and 10° C./min or less. The heating time is preferably 0.5 hours or more, particularly 1 hour or more, and 50 hours or less, particularly 20 hours or less.

The cooling rate after heating is not particularly limited, but is preferably 1°C/min or more, particularly 5°C/min or more, and 100°C/min or less, particularly 50°C/min or less, until the upper limit of the above range is reached. more preferred. If the cooling rate is less than the above range, the (R', HR)--Fe(Co)-- _M.sub.1 phase segregates at the grain boundary triple point, possibly deteriorating the magnetic properties. On the other hand, when the cooling rate exceeds 100° C./min, segregation of the (R′, HR)—Fe(Co)—M _phase can be suppressed during the cooling process, but the squareness of the sintered magnet is particularly affected. It may get worse.

After the high-temperature heat treatment, a low-temperature heat treatment step of heating to a temperature of 400 ° C. or higher, particularly 430 ° C. or higher and 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. implement. The heat treatment atmosphere is preferably a vacuum or an inert gas atmosphere such as Ar gas.

This heating temperature is lower than the peritectic temperature of the (R', HR)-Fe(Co)-M _phase because the (R', HR)-Fe(Co)-M _phase is formed in the grain boundary phase. It is effective to If the heating temperature is less than 400° C., the reaction rate of forming (R′, HR)--Fe(Co)--M ₁ may become 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 grains The boundary phase may segregate significantly at the grain boundary triple point, degrading the magnetic properties.

The rate of temperature increase in 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 raising the temperature at the low-temperature heat treatment temperature is preferably 0.5 hours or more, particularly 1 hour or more, and 50 hours or less, particularly 20 hours or less. The cooling rate after heating is not particularly limited, but is preferably 1°C/min or more, particularly 5°C/min or more, and 100°C/min or less, particularly 80°C/min or less, until the upper limit of the above range is reached. In particular, 50° C./min or less is more preferable. After the heat treatment, the sintered body is usually cooled to room temperature.

EXAMPLES The present invention will be specifically described below 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 metals (Nd or didymium (mixture of Nd and Pr)), electrolytic iron, metals or alloys of Co, M1 and _M2 elements, and Fe _— B alloys (ferroboron) are used to achieve the desired composition An alloy ribbon was produced by weighing as follows, melting in a high-frequency induction furnace in an Ar gas atmosphere, 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 alloy ribbon thus produced was subjected to a hydrogen absorption treatment at normal temperature, and then heated in a vacuum at 600° C. for dehydrogenation to pulverize the alloy. 0.07% by mass of stearic acid as a lubricant was added to the obtained coarse powder and mixed. Next, a mixture of the coarse powder and the lubricant was pulverized by a jet mill in a nitrogen stream to prepare a fine powder having an average particle size of about 3 μm.

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

[Examples 1 to 6, Comparative Examples 1 to 3]
After processing the sintered body obtained in Reference Example 1 into a rectangular parallelepiped shape with a size of 20 mm × 20 mm × 2.2 mm, terbium oxide particles with an average particle diameter of 0.5 µm were added to a slurry mixed with ethanol at a mass fraction of 50%. It was immersed 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 cooled to 200 ° C. at the cooling rate shown in Table 2. High temperature heat treatment was performed. Then, after heating at the holding temperature shown in Table 2 for 2 hours, low-temperature heat treatment was performed by cooling to 200° C. at the cooling rate shown in Table 2 to obtain a sintered magnet. Table 1 shows the composition of the obtained sintered magnet, and Table 2 shows its magnetic properties. The magnetic properties were evaluated by cutting the central portion of the obtained sintered magnet into a rectangular parallelepiped shape of 6 mm×6 mm×2 mm.

1 shows the elemental distribution of Nd (A) and the elemental distribution of Tb (B ) is an image. It can be seen that Tb diffuses through the grain boundary phase and the HR-rich phase is unevenly 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 exists at intergranular grain boundaries and grain boundary triple junctions. It was thickly present in the part of On the other hand, it was confirmed that the grain boundary phase includes (R', HR)-Fe(Co)-M ₁ phase and (R', HR) rich phase, and (R', HR) It was confirmed that the oxide phase was mainly segregated at the grain boundary triple point.

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

In the image of the elemental distribution of Tb, the HR-rich phase, the (R', HR) rich phase and the (R', HR) oxide phase are not clearly distinguished from the elemental distribution of Tb, but the elemental distribution of Nd Focusing on the image, the Nd content is higher in the (R', HR)-rich phase and (R', HR) oxide phase and lower in the HR-rich phase than in the center of the main phase. A distinction is possible. In the cross sections of the R—Fe—B based sintered magnets of the examples and comparative examples, the portion where the Nd content is 80% or less of the Nd content in the center of the main phase is defined as the HR rich phase, and the partial area is Table 2 shows the results of evaluation of the ratio of to the area of the entire main phase. The ratio of the HR-rich phase in the sintered magnets of the examples was higher than that of the sintered magnets of the comparative examples.

[Examples 7 to 9, Comparative Example 4]
After processing the sintered body obtained in Reference Example 2 into a rectangular parallelepiped shape with a size of 20 mm × 20 mm × 2.2 mm, terbium oxide particles with an average particle diameter of 0.5 µm were mixed with ethanol at a mass fraction of 50% in a slurry. It was immersed 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 cooled to 200 ° C. at the cooling rate shown in Table 2. High temperature heat treatment was performed. Then, after heating at the holding temperature shown in Table 2 for 2 hours, low-temperature heat treatment was performed by cooling to 200° C. at the cooling rate shown in Table 2 to obtain a sintered magnet. Table 1 shows the composition of the obtained sintered magnet, and Table 2 shows its magnetic properties. The magnetic properties were evaluated by cutting the central portion of the obtained sintered magnet into a rectangular parallelepiped shape of 6 mm×6 mm×2 mm. Table 2 shows the ratio of the HR-rich phase evaluated by the method described above. The ratio of the HR-rich phase in the sintered magnets of the examples was higher than that of the sintered magnets of the comparative examples.

[Example 10, Comparative Example 5]
After processing the sintered body obtained in Reference Example 1 into a rectangular parallelepiped shape with a size of 20 mm × 20 mm × 2.2 mm, dysprosium oxide particles with an average particle diameter of 0.5 µm were mixed with ethanol at a mass fraction of 50% in a slurry. It was immersed 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 cooled to 200 ° C. at the cooling rate shown in Table 2. High temperature heat treatment was performed. Then, after heating at the holding temperature shown in Table 2 for 2 hours, low-temperature heat treatment was performed by cooling to 200° C. at the cooling rate shown in Table 2 to obtain a sintered magnet. Table 1 shows the composition of the obtained sintered magnet, and Table 2 shows its magnetic properties. The magnetic properties were evaluated by cutting the central portion of the obtained sintered magnet into a rectangular parallelepiped shape of 6 mm×6 mm×2 mm. Table 2 shows the ratio of the HR-rich phase evaluated by the method described above. The ratio of the HR-rich phase in the sintered magnets of the examples was higher than that of the sintered magnets of the comparative examples.

Claims (8)

12 to 17 atomic % R (R is one or more elements selected from rare earth elements including Y, and Nd is essential), 0.1 to 3 atomic % 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 ~ 5.9 + 2 × m) atomic % (m is the content of the element represented by M ₂ (atomic %)) B, 10 atomic % or less Co, 0.5 atomic % or less C, 1.5 atomic % or less O , 0.5 atomic percent or less of _N , and the balance of _Fe . - Fe-B based sintered magnet,
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 formed between the main phase and the crystal grains of the main phase (R', HR)-Fe(Co)-M _phase having a composition (R' contains Y and is one or more elements selected from rare earth elements excluding Dy, Tb and Ho, and Nd is essential, and HR has a grain boundary phase containing one or more elements selected from Dy, Tb and Ho),
(R', HR)-Fe(Co)-M _{1 in the M 1 phase} is
0.5 to 50 atomic % Si, the balance 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 % of Ga and the balance selected from 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 , is one or more elements selected from Pb and Bi,
The main phase contains an HR-rich phase represented by (R′, HR) ₂ (Fe, (Co)) ₁₄ B on its surface, and the content of HR in the HR-rich phase is higher than the content of HR in the core of the phase,
R--Fe-- characterized in that the area of the HR-rich phase evaluated in a cross section within 200 μm from the surface of the R--Fe--B based sintered magnet is 2% or more with respect to the area of the entire main phase. B-based sintered magnet. 2. The R--Fe--B system sintered magnet according to claim 1, wherein said HR-rich phase is non-uniformly formed on the surface of said main phase. 3. The R—Fe—B system firing according to claim 1, wherein the Nd content in the HR-rich phase is 0.8 times or less the Nd content in the center of the main phase. binding magnets. 4. The R--Fe--B system sintered magnet according to any one of claims 1 to 3, characterized in that the total content of Dy, Tb and Ho in the composition of the magnet as a whole is 1 atomic % or less. 5. The magnet structure according to any one of claims 1 to 4, wherein the (R', HR)-Fe(Co)-M ₁ phase is present in a volume fraction of 1% or more and 20% or less. R—Fe—B system sintered magnet. M ₁ in the (R', HR)-Fe(Co)-M ₁ phase is
0.5 to 50 atomic % Si, the balance 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 % of Ga and the balance of Si, Al, Mn, Ni, Cu, Zn, Ge, Pd, Ag, Cd, In, Sn, Sb, Pt, Au, Hg 6. The R--Fe--B system sintered magnet according to any one of claims 1 to 5, wherein the element is one or more elements selected from , Pb and Bi. The R- according to any one of claims 1 to 6, wherein 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. Fe—B system sintered magnet. 8. The R--Fe-- according to any one of claims 1 to 7 , wherein said M ₂ is one or more elements selected from V, Cr, Zr, Nb, Mo, Hf, Ta and W. B-based sintered magnet.

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