CN1860248A - Raw material alloy for r-t-b permanent magnet and r-t-b permanent magnet - Google Patents

Abstract

A raw material alloy for R-T-B permanent magnet of thin-sheet form comprising an R2T14B columnar crystal and an R enriched phase (R: at least one rare earth element including Y, T: Fe or Fe and at least one transition metal element other than Fe, and B: boron or boron and carbon), wherein in the alloy structure observed on any arbitrary section including the normal direction of the thin sheet, the aspect ratio is 10 or higher and the area ratio of R enriched phase whose longitudinal direction is 90+-30 DEG against the surface of the thin sheet is 30% or higher based on all the R enriched phases existing in the alloy.

Description

Raw material alloys for R-T-B permanent magnets and R-T-B permanent magnets

technical field

The invention relates to a raw material alloy for R-T-B series permanent magnets, in particular to a raw material alloy sheet for R-T-B series permanent magnets produced by a strip casting method. In addition, the present invention relates to an R-T-B permanent magnet produced from the above-mentioned raw material alloy for R-T-B permanent magnet.

Background technique

Among permanent magnets, R-T-B permanent magnets with the largest magnetic energy product are used in HD (hard disk), MRI (magnetic resonance imaging), various motors, etc. due to their high characteristics. In recent years, in addition to the improvement of its heat resistance, the market demand for energy saving is also increasing, and the ratio of motor applications including automobiles is increasing.

Here, "R" in the "R-T-B permanent magnet" mainly refers to replacing a part of Nd with other rare earth elements such as Pr and Dy, and is at least one of the rare earth elements including Y. "T" refers to replacing a part of Fe with Co, Ni and other transition metals. "B" is boron, including replacement of part thereof with C or N. "R-T-B permanent magnets" are mainly composed of Nd, Fe, and B, so they are collectively called "Nd-Fe-B magnets" or "R-Fe-B magnets".

The "R-T-B permanent magnet" in this specification includes adding one of elements such as Cu, Al, Ti, V, Cr, Ga, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr, and/or Hf. A magnet of various kinds or a magnet with a combination of the above-mentioned elements. It is known that various properties such as magnetic properties can be improved by adding such elements.

The RTB-based alloy is an alloy in which the R ₂ T ₁₄ B phase, which is a ferromagnetic phase contributing to magnetization, is the main phase, and a non-magnetic, low-melting R-rich phase in which rare earth elements are concentrated coexists. RTB alloys are active metals, so they are usually melted and cast in vacuum or inert gas. In addition, in order to produce sintered magnets from cast RTB-based alloy ingots by powder metallurgy, the alloy ingots are pulverized to about 3 μm (FSSS: measured with a Fischer particle sieve) to form alloy powders, which are then press-molded in a magnetic field. The powder compact obtained by press molding is sintered at a high temperature of about 1000 to 1100° C. using a sintering furnace. The sintered body produced in this way is usually subjected to heat treatment, machining, and plating to improve corrosion resistance as necessary.

The R-rich phase in the R-T-B system sintered magnet has the following important functions.

1) The R-rich phase has a low melting point and becomes a liquid phase during sintering, which contributes to the high density of the magnet and the improvement of magnetization.

2) The unevenness of the grain boundary does not occur, and the nucleation points in the antimagnetic region are reduced to increase the coercive force.

3) Make the main phase magnetically insulated to increase the coercive force.

Therefore, if the dispersion state of the R-rich phase in the molded magnet deteriorates, local poor sintering and magnetic degradation will result. Therefore, it is important to uniformly disperse the R-rich phase in the molded magnet. The distribution of the R-rich phase in the R-T-B-based sintered magnet greatly affects the structure of the R-T-B-based alloy as a raw material.

As a casting method for casting R-T-B alloys, a strip casting method (hereinafter abbreviated as "SC method") has been developed and is being used in actual processes. In the SC method, the alloy melt is cast on an internal water-cooled copper roll, and the alloy melt is rapidly cooled and solidified, thereby casting a thin sheet with a thickness of about 0.1 to 1 mm. By the SC method, since the crystal structure of the alloy is refined, an R-T-B alloy having a structure in which the R-rich phase is finely dispersed can be produced. In this way, the alloy cast by the SC method has finely dispersed R-rich phases inside, so the dispersibility of the R-rich phases in the magnet after crushing and sintering also becomes good, and the improvement of the magnetic properties of the magnet can be realized (Japanese Patent Application Laid-Open No. 5-222488 Publication No. and Japanese Patent Application Laid-Open No. 5-295490).

Alloy flakes cast by the SC method are also excellent in microstructure uniformity. The uniformity of the structure can be compared with the crystal particle size or the dispersion state of the R-rich phase. In the alloy flakes made by the SC method, chilled fine grains (equiaxed grains) sometimes occur on the casting roll side (later, as the casting surface side) of the alloy flakes, but as a whole, the alloy flakes formed during rapid cooling and solidification can be obtained. Moderately fine and uniform organization.

As mentioned above, the R-T-B alloy cast by the SC method has a finely dispersed R-rich phase and excellent structure uniformity. Therefore, when producing a sintered magnet, the uniformity of the R-rich phase in the final magnet is also excellent. Improvement can improve the magnetic properties. In this way, the R-T-B alloy block cast by the SC method has an excellent structure for making sintered magnets. However, as the characteristics of magnets improve, a high degree of control over the structure of the raw material alloy, especially the state of the R-rich phase, is increasingly required.

Previously, the inventors of the present invention studied the relationship between the microstructure of the cast R-T-B alloy and the behavior of hydrogen crushing or pulverization, and found that in order to control the particle size of the alloy powder for sintered magnets to be uniform, the amount of the R-rich phase should be controlled. The state of dispersion is important (JP-A-2003-188006). In addition, it has been found that in the dispersion state of the R-rich phase generated on the mold surface side in the alloy, the extremely fine region (fine R-rich phase region) is easily pulverized, and the pulverization stability of the alloy is reduced, and the powder is reduced. The particle size distribution is broad, and it was recognized that it is necessary to reduce the fine R-phase-rich region in order to improve the magnet properties.

The alloy disclosed in Japanese Unexamined Patent Application Publication No. 2003-188006 can improve crushing stability and magnetic properties. However, merely reducing the fine R-rich phase region cannot fully exert the original function of the above-mentioned R-rich phase, and it is desired to increase the magnetism of the permanent magnet by further controlling the alloy structure.

Contents of the invention

It is an object of the present invention to provide a raw material alloy for R-T-B permanent magnets which can control the R-rich phase existing in the alloy on a finer scale, thereby improving magnetic properties.

The inventors of the present invention observed the R-rich phase existing in the R-T-B alloy on a finer scale, and found that there is a large relationship between the shape of the R-rich phase and the magnetic properties. That is, the present invention is as follows.

(1) A raw material alloy for RTB system permanent magnets, which is a thin plate-shaped raw material alloy for RTB system permanent magnets containing R ₂ T ₁₄ B columnar crystals and R-rich phases (R is a rare earth element including Y) At least one, T is at least one of Fe or Fe and transition metal elements other than Fe, B is boron or boron and carbon), characterized in that, in the alloy structure observed in any section including the normal direction of the thin plate The area ratio of the R-rich phase with an aspect ratio of 10 or more and whose major axis direction is 90±30° with respect to the sheet surface is 30% or more of all R-rich phases present in the alloy.

(2) The raw material alloy for R-T-B permanent magnets according to (1) above, wherein the area ratio of the R-rich phase having an aspect ratio of 10 or more and a long axis direction of 90±30° with respect to the sheet surface It is more than 50% of all R-rich phases present in the alloy.

(3) The raw material alloy for an R-T-B-based permanent magnet as described in (1) above, wherein the area ratio of the R-rich phase having an aspect ratio of 10 or more and a long-axis direction of 90±30° with respect to the sheet surface It is more than 70% of all R-rich phases present in the alloy.

(4) The raw material alloy for an R-T-B permanent magnet as described in (1) to (3) above, wherein the aspect ratio is 20 or more.

(5) A raw material alloy for RTB system permanent magnets, which is a thin plate-shaped raw material alloy for RTB system permanent magnets containing R ₂ T ₁₄ B columnar crystals and R-rich phases (R is a rare earth element including Y) At least one, T is at least one of Fe or Fe and transition metal elements other than Fe, B is boron or boron and carbon), characterized in that, in the alloy structure observed in any section including the normal direction of the thin plate The area ratio of the R-rich phase with an aspect ratio of 10 or more and whose major axis direction is 30° or less or 150° or more relative to the sheet surface is 50% or less of all R-rich phases present in the alloy.

(6) The raw material alloy for R-T-B permanent magnets according to (5) above, characterized in that the R-rich phase has an aspect ratio of 10 or more and a long axis direction of 30° or less or 150° or more with respect to the sheet surface The area ratio of is 30% or less of all R-rich phases present in the alloy.

(7) A raw material alloy for RTB system permanent magnets, which is a thin plate-shaped raw material alloy for RTB system permanent magnets containing R ₂ T ₁₄ B columnar crystals and R-rich phases (R is a rare earth element including Y) At least one, T is at least one of Fe or Fe and transition metal elements other than Fe, B is boron or boron and carbon), characterized in that, in the alloy structure observed in any section including the normal direction of the thin plate The area ratio of the R-rich phase whose aspect ratio is 10 or more and whose long axis direction is 90±30° with respect to the sheet surface is 30% or more of all the R-rich phases present in the alloy, and the aspect ratio is 10 The area ratio of the above R-rich phase whose major axis direction is 30° or less or 150° or more with respect to the sheet surface is 50% or less of all R-rich phases present in the alloy.

(8) The raw material alloy for an R-T-B permanent magnet as described in (7) above, wherein the area ratio of the R-rich phase having an aspect ratio of 10 or more and a long axis direction of 90±30° with respect to the sheet surface is more than 50% of all the R-rich phases present in the alloy, and the area ratio of the R-rich phase with an aspect ratio of 10 or more and whose long axis direction is 30° or less or 150° or more relative to the sheet surface is the largest in the alloy. Less than 30% of the total R-rich phase present.

(9) The raw material alloy for an R-T-B permanent magnet as described in (1) to (8) above, which is produced by a strip casting method.

(10) The raw material alloy for an R-T-B permanent magnet according to (9) above, wherein the average thickness is 0.10 mm or more and 0.50 mm.

(11) An R-T-B-based permanent magnet made of the R-T-B-based alloy described in (1) to (10) above.

Description of drawings

FIG. 1 is a view showing a cross-sectional structure of a rare earth magnet alloy flake containing an aggregated R-rich phase produced by a conventional SC method.

Fig. 2 is a view showing a cross-sectional structure of a rare earth magnet alloy flake containing an R-rich phase with high-order dendritic branches produced by a conventional SC method.

Fig. 3 is a view showing a cross-sectional structure of an alloy flake for a rare earth magnet according to the present invention.

Fig. 4 is a schematic diagram of a casting apparatus for a strip casting method.

Detailed ways

Hereinafter, the raw material alloy for R-T-B permanent magnets of the present invention will be described with reference to the drawings.

implementation.

First, refer to FIG. 1 and FIG. 2 . These figures are reflection electron images obtained by observing the cross-section of a Nd-Fe-B-based alloy (Nd31.5% by mass) cast by the conventional SC method with a SEM (scanning electron microscope). Meanwhile, the left side of the figure is the cast surface side of the alloy, and the right side is the free surface side of the alloy. When performing rapid cooling and solidification of the molten alloy by the SC method, the molten alloy is rapidly cooled from the mold surface side and crystallized.

The white part in FIG. 1 represents an Nd-rich phase (because R becomes Nd, the R-rich phase is sometimes called "Nd-rich phase"). As is clear from FIG. 1 , the Nd-rich phase is stretched and aggregated. On the other hand, in FIG. 2 , a very fine Nd-rich phase exists in the form of dendrites.

In order to produce a sintered magnet from an R-T-B alloy, it is necessary to pulverize the R-T-B alloy and then press it to produce a molded body. As a method of pulverizing the R-T-B-based alloy, it is preferable to finely pulverize the R-T-B-based alloy first by embrittlement of the R-T-B-based alloy by storing hydrogen. The R-T-B-based alloy is coarsely pulverized (broken) by embrittlement caused by hydrogen storage. In the hydrogen crushing step of this R-T-B alloy, hydrogen is absorbed by the R-rich phase, which expands and becomes a brittle hydride. Therefore, in the hydrogen crushing, fine cracks are introduced into the alloy along the R-rich phase or starting from the R-rich phase. In the subsequent pulverization step, the alloy is broken starting from a large number of microcracks generated in the hydrogen crushing, so the dispersion state of the R-rich phase tends to affect the pulverization efficiency and the shape of the fine powder. Therefore, the present inventors observed R-rich phases on a finer scale and found that there is a relationship between the shape of each R-rich phase, the fine cracks formed in hydrogen fragmentation, and magnetic properties.

As can be seen from the elongated R-rich phase shown in FIG. 1 , microcracks are radially formed at the time of gas crushing, and embrittlement itself occurs. Therefore, when pulverized by the jet mill, most of the embrittled elongated R-rich phase is separated from the main phase and pulverized into very fine particles. The extremely fine powder composed of such an R-rich phase has a high rate of being unrecoverable after being separated by a cyclone classifier, and thus becomes a cause of composition variation during pulverization. In addition, the fine powder composed of the R-rich phase is very active, so it becomes the cause of the decrease of the magnetic properties due to the increase of the oxygen concentration, and the strengthening of the safety measures of the process is also required, resulting in the decrease of the production efficiency and the increase of the cost.

On the other hand, the very fine R-rich phase exists in the form of dendrites as shown in FIG. 2 , and the distance between adjacent fine dendritic R-rich phases is also smaller than the pulverized particle size for general sintered magnets. Therefore, in the fine powder pulverized by the jet mill, the ratio of the fine dendritic R-rich phase entering becomes high. As mentioned above, the R-rich phase becomes a liquid phase during sintering and contributes to sintering. For this reason, an R-rich phase exists on the surface of each fine powder, and it is necessary to wet the fine powders to each other during sintering. However, such an effect cannot be expected for the R-rich phase incorporated into the powder, and even if it bleeds out on the surface, a sufficient effect cannot be brought about, causing a decrease in the sintered density of the magnet.

In addition, if a large amount of very fine R-rich phases in the form of dendrites exist, it means that a large number of branches of dendrites exist in the powder, and R ₂ T ₁₄ B phases with different anisotropic orientations coexist in the powder. The ratio of becomes high, so there arises a problem that the degree of orientation of the obtained permanent magnet decreases.

Next, the Nd-Fe-B alloy (Nd31.5% by mass) of the present invention cast by the SC method will be described with reference to FIG. 3 . Fig. 3 shows a reflection electron image obtained by observing a cross-section of a cast piece of the Nd-Fe-B-based alloy of the present invention with a SEM (scanning electron microscope).

As can be clearly seen from FIG. 3, among the Nd-rich phases appearing on the cross-sectional photograph, the layered (flaky) Nd-rich phase elongated within the angle range defined by keeping the thickness direction at the center Proportion is dominant. Only the elongated R-rich phase shown in FIG. 1 or the branched R-rich phase shown in FIG. 2 exists, but its presence ratio is small. Alloys having such a structure can be pulverized by a jet mill after crushing by absorbing hydrogen, which can solve the problems of the alloys having the structures shown in Fig. 1 and Fig. The decrease of magnetic properties, the decrease of sintered density, and the decrease of orientation degree caused by the increase of As a result, an optimal raw material alloy for an R-T-B permanent magnet capable of fully exerting the original function of the R-rich phase can be obtained, and by using such a raw material alloy, an R-T-B permanent magnet having high magnetic properties can be obtained.

In the conventional R-T-B alloys, the structure shown in Fig. 3 partially exists. In addition, it is described in JP-A-09-170055 and JP-A-10-36949 that the dispersed state of the R-rich phase in the R-T-B alloy can be controlled by the cooling rate after the melt solidifies during casting, or heat treatment to control. However, even if the structure as shown in FIG. 3 partially exists in the existing R-T-B alloy, the structure as shown in FIGS. 1 and 2 accounts for most of the structure, so the R-rich phase as described above cannot be fully exerted. original role.

Hereinafter, the raw material alloy for R-T-B permanent magnets of the present invention will be described in detail.

(1) Strip casting method

First, casting of a raw material alloy for an R-T-B permanent magnet by a strip casting method will be described with reference to FIG. 4 . Figure 4 shows a schematic view of an apparatus for casting by strip casting.

Generally, R-T-B alloys are melted in a refractory crucible 1 in a vacuum or an inert gas atmosphere because of their active properties. The melted solution is kept at 1300-1500° C. for a predetermined time, and then, if necessary, passes through a rectification mechanism and a tundish 2 equipped with a dross removal mechanism, and is supplied to a casting rotary roll 3 whose interior is water-cooled.

The supply speed of the molten metal and the rotation speed of the rotating roll are appropriately controlled according to the desired thickness of the alloy. The rotational peripheral speed of the rotating roller is preferably set to about 0.5 to 3 m/s. As the material of the rotating roll for casting, copper or a copper alloy is suitable because it has good thermal conductivity and is easy to obtain. Due to the difference in the material of the rotating roll or the surface state of the roll, the surface of the casting rotating roll is prone to metal adhesion. Therefore, if necessary, if a cleaning device is installed, the quality of the cast R-T-B alloy can be stabilized. The alloy 4 solidified on the rotating roll is detached from the roll on the opposite side of the tundish and recovered in the recovery container 5 . The structure state of the R-rich phase can be controlled by providing a heating and cooling mechanism in the recovery container.

When producing the alloy of the present invention, it is necessary to appropriately set cooling on the casting roll (referred to as "primary cooling") and cooling in a recovery container (referred to as "secondary cooling").

The primary cooling, specifically, is to make the alloy temperature reach 600-850°C when it leaves the casting roll. It is desirable for the alloy to leave the casting rolls at a temperature higher than the melting point of the R-rich phase. Due to the difference in composition, the melting point of the R-rich phase is higher or lower, but it is above 600°C. When the alloy temperature when leaving the casting roll is lower than the melting point of the R-rich phase, the solidification of the R-rich phase has ended, so the structure shown in Figure 2 is formed. On the other hand, when the temperature is higher than 850° C., the R-rich phase aggregates into a stretched state after leaving the roll, forming the structure shown in FIG. 1 . A more preferred temperature range for the alloy as it leaves the casting rolls is 600-800°C. The most preferred temperature range is 640-750°C. However, depending on the alloy composition, the preferred temperature range is higher or lower.

The dispersion state or shape of the R-rich phase greatly depends on TRE (total rare earth content). For example, in an alloy with low TRE and few R-rich phases, the heat collected on the casting roll is small, so the temperature of the alloy when the casting roll is detached tends to be high, and the R-rich phase tends to aggregate and form a stretched shape. On the other hand, in an alloy with a high TRE and a large amount of R-rich phases, the heat collected on the casting roll is large, so the tendency to form a structure having high-order dendritic branches is strong. Therefore, in order to make the temperature of the alloy when it leaves the casting roll fall within the above-mentioned suitable temperature range, the thickness of the alloy needs to be thinned when there is little TRE, and the thickness of the alloy needs to be thickened when there is a lot of TRE. Specifically, when the target TRE is 30% by weight or less, the degree of primary cooling is increased, so it is preferable to make the average thickness of the alloy 0.10 to 0.30 mm. More preferably, it is 0.15 to 0.27 mm. Most preferably, it is 0.20 to 0.25 mm. With a TRE target of 30% to 33% by weight, the preferred average thickness of the alloy is 0.25-0.35 mm. More preferably, it is 0.26 to 0.32 mm. The preferred average thickness of the alloy is 0.28 to 0.50 mm when the TRE target is 33% by weight or more. More preferably, it is 0.28 to 0.35 mm.

In addition, the degree of primary cooling can also be controlled by properly selecting the surface roughness of the casting roll and controlling the heat collected from the alloy by the casting roll. This method is particularly effective when the TRE target is 30% by weight or less or 33% by weight or more. By making the surface roughness of the casting rolls large, it is possible to suppress the collection of heat more than necessary. When TRE is 33% by weight or more, high heat transfer to the casting roll due to a large amount of R-rich phase can be moderately suppressed by increasing the surface roughness of the casting roll. The surface roughness of the target in this case is 20 micrometers or more in ten-point average roughness Rz. When TRE is 30% by weight or less, on the contrary, in order not to hinder the primary cooling on the roll, the surface roughness is set to be about 20 μm or less, and it is preferable to prevent transitional aggregation of the R-rich phase. However, the surface roughness is also affected by other factors such as the surface material of the roll, and therefore is not limited to the above-mentioned numerical values.

Casting roll surface temperature affects the wettability of alloy melt and roll. If the temperature is too low, the wetting of the molten alloy and the roll will deteriorate, and there will be a tendency to cause macroscopic non-uniformity in the contact between the two. As a result, a temperature distribution occurs in the alloy, which becomes a fluctuation factor of the alloy temperature at the time of detachment from the above-mentioned preferred casting roll, and the formation of the R-rich phase having the specific shape of the present invention becomes difficult. On the other hand, if the temperature is too high, the wettability of the alloy melt and the roll becomes good, and partial thermal adhesion may occur. Thermal adhesion of the alloy to the surface of the roll leads to changes in its partial heat conduction and wettability, and becomes a major factor in the change of the alloy structure. Therefore, it is still difficult to form the R-rich phase having the specific shape of the present invention. In addition, if the amount of alloy to be thermally adhered increases, stable operation becomes difficult, resulting in a decrease in productivity. Therefore, the casting roll surface temperature is suitably 50 to 400°C, preferably 100 to 300°C, and most preferably 150 to 200°C. The temperature of the roll surface shown here is the temperature of the part where the melt touches the roll. Although it is difficult to measure it directly, it can be obtained by embedding a thermocouple directly under the casting surface or by contacting the lower part of the tundish without directly contacting the alloy or the melt. The measurement value of the thermocouple on the roller surface of the liquid part can be obtained.

On the other hand, for secondary cooling, for example, a partition is installed in the recovery container, the interval between the partitions is appropriately set, and the inside of the partition is cooled with an inert gas such as Ar or water to control the cooling of the recovered alloy. Speed works. When producing the alloy of the present invention, when the alloy temperature when recovered in the recovery container is 650-700°C, the preferred cooling rate to reach 600°C is 3-30°C/min, preferably 3-20°C/min . When the temperature of the alloy when recovered in the recovery container is 700 to 800°C, the cooling rate to 600°C is preferably 10 to 40°C/min, preferably 10 to 30°C/min. When the temperature of the alloy when recovered in the recovery container is 800 to 850°C, the cooling rate to reach 600°C is preferably 20 to 50°C/min, preferably 30 to 50°C/min. In these temperature ranges, if the upper limit is exceeded, the structure shown in FIG. 2 is likely to be formed. Also, if it is below the lower limit, the structure shown in FIG. 1 is likely to be formed.

In addition, the alloy of the present invention has a defined structure, and the production method is not limited to the above-mentioned method.

The thickness of the alloy flakes of the present invention is preferably not less than 0.1 mm and not more than 0.5 mm. If the alloy flakes are thinner than 0.1 mm, the solidification rate increases excessively, and the dispersion of the R-rich phase becomes too fine. In addition, when the thickness of the alloy flakes is thicker than 0.5 mm, problems such as a decrease in the dispersibility of the R-rich phase due to a decrease in the solidification rate are caused.

(2) R-rich phase in the alloy

The present invention is a raw material alloy for RTB-based permanent magnets containing R ₂ T ₁₄ B columnar crystals and a thin-plate R-rich phase (R is at least one rare earth element including Y, T is Fe or Fe and Fe and other than Fe) At least one of the transition metal elements, B is boron or boron and carbon). Moreover, in the alloy structure observed at any cross section including the normal direction of the sheet, the area ratio of the R-rich phase with an aspect ratio of 10 or more and whose long axis direction is 90±30° with respect to the sheet surface is the highest in the alloy. More than 30% of all R-rich phases. As a result, there is little change in the composition during pulverization in the sintered magnet manufacturing process, there is no decrease in magnetic properties due to an increase in the concentration of oxygen and nitrogen, and there is no decrease in sintered density and orientation. The most suitable raw material alloy for RTB-based permanent magnets with the original role of the R phase. In addition, by using this raw material alloy, an RTB-based permanent magnet having high magnetic properties can be obtained.

When the aspect ratio of the R-rich phase in the alloy is less than 10, the R-rich phase is agglomerated and stretched, and if the ratio of such an R-rich phase increases, the R-rich phase will fall off during pulverization, resulting in excessive pulverization. Composition changes will increase.

In addition, even if the aspect ratio is 10 or more, the long axis direction is outside the range of 90±30° with respect to the surface of the sheet, and there is a high possibility that the R-rich phase is more than necessary and its intervals are fine branchlets. Such an R-rich phase can explain high-order dendrites in metal histology, but in the actual alloy structure, there is a high possibility of individual differences in the identification of primary dendrites and secondary or higher-order dendrites. , define geometrically, and stipulate this range.

As a result, if the area ratio of the R-rich phase whose aspect ratio is 10 or more and whose major axis direction is 90±30° with respect to the sheet surface is 30% or less of the total R-rich phase present in the alloy, the magnetic The reduction in characteristics becomes noticeable.

Preferably, the area ratio of the R-rich phase having an aspect ratio of 10 or more and whose major axis direction is 90±30° with respect to the sheet surface is 50% or more of all R-rich phases present in the alloy. Most preferably, the area ratio of the R-rich phase having an aspect ratio of 10 or more and whose long axis direction is 90±30° with respect to the sheet surface is 70% or more of all R-rich phases present in the alloy.

More preferably, in the above alloy, the aspect ratio is 20 or more. More preferably, in the above alloy, the aspect ratio is 30 or more.

Alternatively, the area ratio of the R-rich phase having an aspect ratio of 10 or more and whose major axis direction is 30° or less or 150° or more relative to the sheet surface is 50% or less of all R-rich phases present in the alloy. Even if the R-rich phase has an aspect ratio of 10 or more and a major axis direction of 30° or less or 150° or more with respect to the surface of the sheet, it is highly likely to be high-order dendrites whose intervals are finely branched.

Preferably, the area ratio of R-rich phases having an aspect ratio of 10 or more and whose major axis direction is 30° or less or 150° or more relative to the sheet surface is 30% or less of all R-rich phases present in the alloy.

Alternatively, the area ratio of the R-rich phase whose aspect ratio is 10 or more and whose major axis direction is 90±30° with respect to the sheet surface is 30% or more of all the R-rich phases present in the alloy, and the aspect ratio is The area ratio of the R-rich phase whose major axis direction is 30° or less or 150° or more with respect to the sheet surface is 10 or more and is 50% or less of all R-rich phases present in the alloy.

Preferably, the area ratio of the R-rich phase with an aspect ratio of 10 or more and whose major axis direction is 90±30° with respect to the sheet surface is 50% or more of all R-rich phases present in the alloy, and the length, width The area ratio of the R-rich phase having a ratio of 10 or more and whose major axis direction is 30° or less or 150° or more relative to the metal surface is 30% or less of all R-rich phases present in the alloy.

The above-mentioned R-rich phase with an aspect ratio of 10 or more, or an aspect ratio of 20 or more, or an aspect ratio of 30 or more has a major axis dimension of 5% or more of the thickness dimension of the sheet, and preferably has a length of 10% or more.

The aspect ratio of the R-rich phase in the alloy, the angle of the major axis direction relative to the surface of the sheet, and the area ratio of such an R-rich phase. Since the R-rich phase is brighter than the main phase on BEI, it can be determined using an image analysis device. Analysis is performed after identifying the main phase and the R-rich phase. For example, photograph the BEI of 10 randomly selected alloy flake sections at an appropriate magnification, and in each of the 10 photographs, the sum of the total area of the R-rich phase in the photograph is the specified aspect ratio and its long axis The total area of the R-rich phase whose direction is within a predetermined angle range is subjected to image analysis processing and measured. Then, the total value of the total area of the R-rich phase in the 10 photographed photographs is divided by the total value of the R-rich phase obtained in each photograph, which has a predetermined aspect ratio and whose long axis direction is within a predetermined angle range. The value obtained from the total value of the total area of the phases can be regarded as the predetermined area ratio of the R-rich phase.

(3) R ₂ T ₁₄ B phase in the alloy

The thin plate-shaped raw material alloy for RTB-based permanent magnets of the present invention has the R ₂ T ₁₄ B phase as the main phase as the ferromagnetic phase. The R ₂ T ₁₄ B phase is columnar, and the R ₂ T ₁₄ B columnar crystals preferably have a major axis within an angle of 90±30° with respect to the sheet surface. In addition, the length of the major axis is 30% or more, preferably 50% or more, of the thickness dimension of the thin plate. Furthermore, the above-mentioned preferred R ₂ T ₁₄ B columnar crystals contain 30% or more of the R ₂ T ₁₄ B columnar crystals in the entire sheet, preferably 50% or more.

In addition, the R ₂ T ₁₄ B columnar crystal in this case refers to a block in which the crystal orientation is consistent when observed with a polarizing microscope utilizing the magnetic Kerr effect.

Hereinafter, examples and comparative examples of the present invention will be described.

(Example 1)

In order to make the alloy composition Nd: 31.5% by mass, B: 1.00% by mass, Co: 1.0% by mass, Al: 0.30% by mass, Cu: 0.10% by mass, and iron as the balance, the alloying metals are neodymium, boron-iron alloy, cobalt, and aluminum , copper and iron raw materials are melted in a high-frequency melting furnace, and the melt is cast using a strip casting method to make alloy flakes.

The diameter of the rotating roll for casting is 300mm, the material is pure copper with a wall thickness of 50mm, and the inside is water-cooled. The peripheral speed of the roll during casting is 1.0m/s, and an alloy flake with an average thickness of 0.27mm is produced. At this time, the average roughness Rz of the casting roll surface was 12 micrometers. When visually observed, the alloy was evenly placed on the casting rolls, and thermal adhesion to the casting rolls was not observed.

In addition, a thermocouple was brought into contact with the bottom of the surface of the casting roll to measure the surface temperature of the casting roll during casting. Furthermore, the amount of cooling water for casting rolls, the temperature difference between the inlet and outlet, and the temperature of the water discharged from the casting rolls were also measured. From these measured values, the surface temperature of the casting rolls at the position where the molten metal in the tundish contacts the casting rolls was calculated as 170°C.

In addition, when the temperature of the alloy flakes detached from the roll was measured with an infrared thermometer, it was 730°C. A partition plate through which Ar gas for cooling flows is installed in the recovery container for storage. A thermocouple was inserted from the side of the recovery container to the inside, and the temperature change of the alloy was measured. As a result, the highest temperature was 720°C, and the average cooling rate to 600°C was 22°C/min.

Ten of the obtained alloy flakes were embedded and ground, and each alloy flake was photographed by reflection electron imaging (BEI) at a magnification of 100 times with a scanning electron microscope (SEM). The photograph taken was placed on an image analysis device for measurement. As a result, the area ratio of the R-rich phase with an aspect ratio of 10 or more and whose long axis direction is 90 ± 30° relative to the metal surface is the highest in the alloy. 80% of the total R-rich phase. In addition, the area ratio of the R-rich phase with an aspect ratio of 20 or more and whose major axis direction is 90±30° with respect to the metal surface is 65% of all R-rich phases present in the alloy. On the other hand, the area ratio of the R-rich phase with an aspect ratio of 10 or more and whose major axis direction is 30° or less or 150° or more with respect to the metal surface is 6% of all R-rich phases present in the alloy.

(comparative example 1)

Raw materials were blended to have the same composition as in Example 1, melted in the same manner as in Example 1, and cast by the SC method. Among them, the wall thickness of the casting roll is specified as 90mm, and the average roughness Rz of the surface of the casting roll is specified as 7 microns. In addition, the average thickness of the alloy flakes was specified to be 0.35 mm. When observed visually, a portion where the temperature of the alloy on the casting roll was abnormally high occurred in a small amount, and a thermal seizing phenomenon was observed in a portion.

The surface temperature of the casting roll at the position where the molten metal in the tundish contacts the casting roll, which was determined by the same method as in Example 1, was 400°C.

In addition, the temperature of the heat-adhered alloy flakes that did not come off the roll was measured with an infrared thermometer to be 820°C. In addition, a special cooling mechanism is installed in the recovery container for storing alloy flakes detached from the roll. The temperature change of the alloy was measured with a thermocouple inserted from the side of the recovery vessel to the inside. As a result, the maximum temperature was 810°C, and the average cooling rate to 600°C was 6°C/min.

As a result of evaluating the obtained alloy flakes without thermal adhesion in the same manner as in Example 1, many R-rich phases aggregated to form a stretched shape, so the area ratio of the R-rich phases with an aspect ratio of 10 or more was the highest in the alloy. 26% of the total R-rich phase present.

(comparative example 2)

Raw materials were blended to have the same composition as in Example 1, melted in the same manner as in Example 1, and cast by the SC method. Among them, the wall thickness of the casting roll is specified as 25mm, and the average roughness Rz of the surface of the casting roll is specified as 10 microns. In addition, the average thickness of the alloy flakes was specified to be 0.22 mm. When visually observed, a part of the alloy on the casting roll had a slightly higher temperature.

The surface temperature of the casting roll at the position where the molten alloy in the tundish contacts the casting roll, which was determined in the same manner as in Example 1, was 80°C.

In addition, the average temperature of the alloy flakes detached from the roll was measured with an infrared thermometer to be 670°C. In addition, in the recovery container for storing the alloy flakes detached from the roll, a partition plate through which cooling water flows is provided. The temperature change of the alloy was measured by inserting a thermocouple from the side of the recovery vessel to the inside. As a result, the highest temperature was 660°C, and the average cooling rate to 600°C was 35°C/min.

As a result of evaluating the obtained alloy flakes in the same manner as in Example 1, the R-rich phase contained a large number of branch-shaped higher-order dendrites, the aspect ratio was 10 or more, and the long-axis direction was 90 ± 30 with respect to the metal surface. The area ratio of the R-rich phase is 23% of the total R-rich phase present in the alloy. On the other hand, the area ratio of the R-rich phase with an aspect ratio of 10 or more and a major axis direction of 30° or less or 150° or more with respect to the metal surface was 54% of all R-rich phases present in the alloy.

Next, examples of sintered magnets will be described.

(Example 2)

The alloy flakes obtained in Example 1 were coarsely pulverized by known hydrogen absorbing crushing treatment, 0.07% by mass of zinc stearate powder was added to the obtained coarsely pulverized powder, and after mixing in a nitrogen atmosphere with a shaker mixer, the A jet mill performs fine pulverization. The atmosphere at the time of pulverization by the jet mill was defined as a nitrogen atmosphere mixed with 10,000 ppm of oxygen. The oxygen concentration of the obtained powder was 5000 ppm. The obtained powder was mixed with cold embedding resin, hardened, and ground, and the cross section of the powder was observed by SEM-BEI to investigate the dispersion state of the R-rich phase in the powder. As a result, the R-rich phase mainly adheres to the surface of the particles composed of the main phase.

Next, the obtained powder was press-molded at a pressure of 1.0 t/cm ² in a magnetic field with an orientation magnetic field of 1.5 T, and the compact was held at 1060° C. for 4 hours for sintering. The sintered density of the obtained sintered body was 7.5 g/cm ³ or more, which was a sufficiently large density. Furthermore, this sintered body was heat-treated at 560° C. for 1 hour in an argon atmosphere to obtain a sintered magnet.

The magnetic properties of the sintered magnets were measured with a BH curve recorder, and the results are shown in Table 1.

(comparative example 3)

The alloy flakes obtained in Comparative Example 1 were pulverized by the same method as in Example 2 to obtain fine powder. At this time, the cross section of the powder was observed by the same method as in Example 2, and it was confirmed that many R-rich phases were separated from the main phase and existed as relatively small particles consisting only of the R-rich phase. Then, through the same molding and sintering steps as in Example 2, a sintered magnet was produced.

The magnetic properties of the sintered magnet produced in Comparative Example 3 were measured with a BH curve recorder, and the results are shown in Table 1.

(comparative example 4)

The alloy flakes obtained in Comparative Example 2 were pulverized by the same method as in Example 2 to obtain fine powder. At this time, the cross section of the powder was observed by the same method as in Example 2, and it was confirmed that the ratio of the particles containing the R-rich phase inside was also about 7 times that of Example 2. Then, through the same molding and sintering steps as in Example 2, a sintered magnet was produced.

The magnetic properties of the sintered magnet produced in Comparative Example 4 were measured with a BH curve recorder, and the results are shown in Table 1.

Table 1 Density (g/cm ³ ) Br(T) iHc(kA·m) (BH) _max (kJ/m ³ ) Amount of TRE in sintered magnet (mass%) Oxygen content in sintered magnet (ppm) Example 2 7.55 1.39 1194 362 31.2 5500 Comparative example 3 7.48 1.37 1098 344 30.6 5400 Comparative example 4 7.52 1.37 1154 352 31.1 5500

As shown in Table 1, Comparative Example 3 has a lower density than Example 2, and has lower magnetization and coercive force in terms of characteristics. The reason can be inferred as: the dispersion of the R-rich phase in the alloy stage deteriorates, so in the pulverization process, the R-rich phase is separated by the cyclone classifier of the pulverizer as an active small powder, and the TRE is easy to decrease, or the segregation of the R-rich phase The sinterability is reduced, so it does not function sufficiently and effectively during sintering. On the other hand, it can be inferred that although Comparative Example 4 is not as good as Comparative Example 3, it also exhibits the same behavior, and that the R-rich phase does not sufficiently contribute to sintering.

Industrial availability

According to the raw material alloy for R-T-B permanent magnets of the present invention, the R-rich phase in the alloy can be effectively utilized to the maximum extent, therefore, the sintered magnet manufactured by this alloy exhibits more excellent magnet properties than the existing sintered magnets . That is, since the R-rich phase fully exerts its original function, in the sintered magnet manufacturing process, there is little composition change during fine pulverization, no decrease in magnetic properties due to increase in oxygen concentration, no decrease in sintered density and no orientation degree It has excellent effects that cannot be obtained with conventional alloys. In addition, by using the above-mentioned raw material alloy, an R-T-B permanent magnet having high magnetic properties can be obtained.

The present invention can be suitably applied to various electronic devices, electric machines, etc. that require high-performance sintered magnets.

Claims (11)

1. one kind contains R ₂T ₁₄B column crystallization and rich R laminal R-T-B series permanent magnet mutually is characterized in that with raw alloy (R is rare earth element at least a that comprises Y, and T is transition metal at least a beyond Fe or Fe and the Fe, and B is boron or boron and carbon):

In the alloy structure of observing on the arbitrary section that comprises described thin plate normal direction, long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of 90 ± 30 ° rich R phase is more than 30% of existing whole rich R phases in the alloy with respect to described thin sheet surface.

2. R-T-B series permanent magnet raw alloy according to claim 1 is characterized in that: long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of 90 ± 30 ° rich R phase is more than 50% of existing whole rich R phases in the alloy with respect to thin sheet surface.

3. R-T-B series permanent magnet raw alloy according to claim 1 is characterized in that: long-width ratio is more than 10 and the relative thin sheet surface of its long axis direction is that the area occupation ratio of 90 ± 30 ° rich R phase is more than 70% of existing whole rich R phases in the alloy.

4. according to each described R-T-B series permanent magnet raw alloy in the claim 1～3, it is characterized in that: long-width ratio is for more than 20.

5. one kind contains R ₂T ₁₄B column crystallization and rich R laminal R-T-B series permanent magnet mutually is characterized in that with raw alloy (R is rare earth element at least a that comprises Y, and T is transition metal at least a beyond Fe or Fe and the Fe, and B is boron or boron and carbon):

In the alloy structure of observing on the arbitrary section that comprises described thin plate normal direction, long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of the rich R phase below 30 ° or more than 150 ° is below 50% of existing whole rich R phases in the alloy with respect to thin sheet surface.

6. R-T-B series permanent magnet raw alloy according to claim 5 is characterized in that: long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of the rich R phase below 30 ° or more than 150 ° is below 30% of existing whole rich R phases in the alloy with respect to thin sheet surface.

7. one kind contains R ₂T ₁₄B column crystallization and rich R laminal R-T-B series permanent magnet mutually is characterized in that with raw alloy (R is rare earth element at least a that comprises Y, and T is transition metal at least a beyond Fe or Fe and the Fe, and B is boron or boron and carbon):

In the alloy structure of on the arbitrary section that comprises described thin plate normal direction, observing, long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of 90 ± 30 ° rich R phase is more than 30% of existing whole rich R phases in the alloy with respect to thin sheet surface, and long-width ratio is more than 10 and its long axis direction is that area occupation ratio for the rich R phase below 30 ° or more than 150 ° is below 50% of existing whole rich R phases in the alloy with respect to thin sheet surface.

8. R-T-B series permanent magnet raw alloy according to claim 7, it is characterized in that: long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of 90 ± 30 ° rich R phase is more than 50% of existing whole rich R phases in the alloy with respect to thin sheet surface, and long-width ratio is more than 10 and its long axis direction is that the area occupation ratio of the rich R phase below 30 ° or more than 150 ° is below 30% of existing whole rich R phases in the alloy with respect to thin sheet surface.

9. according to each described R-T-B series permanent magnet raw alloy in the claim 1～8, it is characterized in that: use the thin strip casting manufactured.

10. R-T-B series permanent magnet raw alloy according to claim 9 is characterized in that: mean thickness is below the above 0.50mm of 0.10mm.

11. R-T-B series permanent magnet of making of raw alloy by each described R-T-B series permanent magnet in the claim 1～10.

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