CN1295713C - R－T－B series rare earth permanent magnet - Google Patents

Abstract

An R-T-B based rare earth permanent magnet comprises a magnet composed of a magnet₂T₁₄A main phase consisting of a B phase (R is 1 or 2 or more kinds of rare earth elements (wherein the rare earth elements are a concept containing Y), T is Fe or 1 or 2 or more kinds of transition metal elements essential to Fe and Co), and a sintered body of a grain boundary phase containing more R than the main phase, wherein R is a rare earth element₂T₁₄Zr-enriched products are present in the B phase. The product enriched in Zr has a flaky or needle-like morphology. Further, the R-T-B based rare earth permanent magnet containing the product can suppress the growth of crystal grains while minimizing the decrease in magnetic properties, and can obtain a wide sintering temperature range.

Description

R-T-B series rare earth permanent magnet

technical field

The present invention relates to R (R is one or more than two kinds of rare earth elements, but the concept of rare earth elements contains Y), T (T is Fe or at least one or more of Fe and Co are necessary Transition metal elements) and R-T-B rare earth permanent magnets mainly composed of B (boron).

Background technique

Among the rare-earth permanent magnets, the R-T-B series rare-earth permanent magnets have excellent magnetic properties, and the main component Nd is abundant and relatively cheap, so the demand is increasing year by year.

In order to improve the magnetic properties of R-T-B series rare earth permanent magnets, research and development are vigorously carried out. For example, Japanese Unexamined Patent Publication No. 1-219143 reports that adding 0.02 to 0.5 atomic % of Cu to an R-T-B-based rare earth permanent magnet can improve magnetic properties and improve heat treatment conditions. However, the method described in Japanese Unexamined Patent Application Publication No. 1-219143 is useful for obtaining high magnetic properties required for high-performance magnets, specifically for obtaining relatively high coercive force (HcJ) and residual magnetic flux density (Br). inadequate.

Here, the magnetic properties of the R-T-B based rare earth permanent magnet obtained by sintering may depend on the sintering temperature. On the other hand, on an industrial scale, it is difficult to make the heating temperature uniform over the entire area in the sintering furnace. Therefore, for R-T-B series rare earth permanent magnets, it is required to obtain desired magnetic properties even if the sintering temperature fluctuates. Here, the sintering temperature range in which desired magnetic properties can be obtained is called the sintering temperature range.

In order to make the R-T-B series rare earth permanent magnet into a higher-performance permanent magnet, it is necessary to reduce the oxygen content in the alloy. However, when the oxygen content in the alloy is reduced, abnormal grain growth tends to occur in the sintering process, and the squareness ratio (also referred to as the squareness ratio) decreases. This is because the oxide formed by the oxygen in the alloy suppresses the grain growth.

Here, a method of adding a new element to a Cu-containing R-T-B-based rare-earth permanent magnet has been studied as a means for improving magnetic properties. It is reported in JP-A-2000-234151 that Zr and/or Cr are added in order to obtain high coercive force and residual magnetic flux density.

Similarly, it is reported in JP-A-2002-75717 that by making the fine ZrB compound, NbB compound or HfB compound (hereinafter referred to as M-B Compounds) are evenly dispersed and precipitated, inhibiting the grain growth in the sintering process, and improving the magnetic properties and sintering temperature range.

According to Japanese Patent Laid-Open No. 2002-75717, the sintering temperature range can be widened by dispersing and precipitating the M-B compound. However, in Example 3-1 disclosed in Japanese Unexamined Patent Publication No. 2002-75717, the sintering temperature range is about 20° C. which is relatively narrow. Therefore, in mass-produced sintering furnaces and the like, it is desired to further increase the sintering temperature range in order to improve magnetic properties. Also, in order to obtain a sufficiently wide sintering temperature range, it is effective to increase the amount of Zr added. However, as the amount of Zr added increases, the remanence magnetic flux density decreases, and the original high characteristics cannot be obtained.

Contents of the invention

Therefore, an object of the present invention is to provide an R-T-B based rare earth permanent magnet capable of minimizing deterioration of magnetic properties, suppressing growth of crystal grains, and further improving a sintering temperature range.

The inventors of the present invention have found that when a Zr-enriched product exists in the R ₂ T ₁₄ B phase constituting the main phase of the RTB-based rare earth permanent magnet, the decrease in magnetic properties can be minimized and the growth of crystal grains can be suppressed. And can improve the sintering temperature range. That is, the present invention provides a kind of RTB series rare earth permanent magnet, which is composed of R ₂ T ₁₄ B phase (R is one or more than two kinds of rare earth elements, wherein the rare earth elements are the concepts containing Y, T It is a sintered body consisting of a main phase composed of at least one transition metal element mainly composed of Fe or Fe and Co) and a grain boundary phase containing more R than the main phase, in which R ₂ T ₁₄ B phase In enriched Zr products.

For the RTB-based rare earth permanent magnet of the present invention, the amount of oxygen contained in the sintered body is preferably 2000 ppm or less. This is because the effects of inhibiting grain growth and expanding the sintering temperature range due to the presence of Zr-enriched products in the R ₂ T ₁₄ B phase are significant at low oxygen content of 2000 ppm or less.

For the R-T-B series rare earth permanent magnet of the present invention, the preferred composition is: R: 28-33% by weight, B: 0.5-1.5% by weight, Al: 0.03-0.3% by weight, Cu: 0.3% by weight or less (excluding 0 ), Zr: 0.05 to 0.2% by weight, Co: 4% by weight or less (excluding 0), and the remainder is substantially composed of Fe.

Furthermore, it is more preferable to contain Zr in the range of 0.1 to 0.15% by weight in the R-T-B based rare earth permanent magnet of the present invention.

Description of drawings

FIG. 1 is a graph showing a combination of a low-R alloy and a high-R alloy used in the first embodiment and the composition of the resulting permanent magnet.

Fig. 2 is a graph showing the magnetic properties of the permanent magnet obtained in the first example.

Fig. 3 is a graph showing the relationship between the amount of the additive element M (Zr or Ti) and the residual magnetic flux density (Br) of the permanent magnet obtained in the first example.

Fig. 4 is a graph showing the relationship between the amount of the added element M (Zr or Ti) and the coercive force (HcJ) of the permanent magnet obtained in the first example.

Fig. 5 is a graph showing the relationship between the amount of the additive element M (Zr or Ti) and the squareness ratio (Hk/HcJ) of the permanent magnet obtained in the first example.

FIG. 6 is a TEM (Transmission Electron Microscope: transmission electron microscope) photograph showing a sample (sample having a Zr amount of 0.10% by weight) of the first example.

Fig. 7(a) is an EDS (Energy Dispersive X-ray Fluorescence Spectroscopy: Energy Dispersive X-ray Fluorescence Spectroscopy: energy dispersive X-ray analyzer spectrometry) showing the product present in the sample of the first example (a sample with a Zr amount of 0.10% by weight). Distribution.

Fig. 7(b) is an EDS distribution diagram showing the R ₂ T ₁₄ B phase of the sample of the first example (a sample having a Zr content of 0.10% by weight).

Fig. 8 is a high-resolution TEM photograph showing a sample of the first example (a sample having a Zr content of 0.10% by weight).

Fig. 9 is a TEM photograph showing a sample (a sample having a Zr content of 0.10% by weight) of the first example.

Fig. 10 is a TEM photograph showing a sample (a sample having a Zr content of 0.10% by weight) of the first example.

Fig. 11(a) is a photograph showing the Zr mapping (mapping) results of EPMA (Electron Probe Micro Analyzer: Electron Probe Micro Analyzer) of the sample of the first embodiment (a sample with a Zr content of 0.10% by weight) (lower section) and a photo of the composition imaged in the same field of view as the Zr mapping result (upper section).

Fig. 11(b) is a photo (lower stage) showing the Zr mapping result of the EPMA of the sample of Comparative Example 2 (a sample with a Zr content of 0.10% by weight) and a photo of the composition image in the same field of view as the Zr mapping result ( Upper section).

Fig. 12 is a graph showing the magnetic properties of the permanent magnet obtained in the second example.

Fig. 13 is a graph showing the relationship between the sintering temperature and the residual magnetic flux density (Br) in the second embodiment.

Fig. 14 is a graph showing the relationship between sintering temperature and coercive force (HcJ) in the second example.

Fig. 15 is a graph showing the relationship between sintering temperature and squareness ratio (Hk/HcJ) in the second embodiment.

Fig. 16 is a graph showing the relationship between the residual magnetic flux density (Br) and the squareness ratio (Hk/HcJ) at each sintering temperature in the second embodiment.

Fig. 17 is a graph showing a combination of a low-R alloy and a high-R alloy used in the third embodiment and the composition of the resulting permanent magnet.

Fig. 18 is a graph showing the magnetic properties of the permanent magnet obtained in the third example.

Fig. 19 is a graph showing combinations of low-R alloys and high-R alloys used in the fourth embodiment and compositions of permanent magnets obtained.

Fig. 20 is a graph showing the magnetic properties of the permanent magnet obtained in the fourth example.

Detailed ways

Hereinafter, embodiments of the present invention will be described.

<organization>

As everyone knows, the permanent magnet obtained according to the present invention contains at least R ₂ T ₁₄ B phase (R is one or more than two kinds of rare earth elements (but the concept of rare earth elements contains Y), T is Fe or A main phase composed of Fe and Co as essential transition metal elements (one or more), and a grain boundary phase containing more R than the main phase. The present invention is characterized in that a Zr-enriched product exists in the R ₂ T ₁₄ B phase. The RTB-based rare-earth permanent magnet containing this product can minimize the decrease in magnetic properties, suppress the growth of crystal grains, and obtain a wide sintering temperature range. It is essential that the product is present in the R ₂ T ₁₄ B phase, but it is not required that the entire product exists in the R ₂ T ₁₄ B phase. In addition, the product may exist in the grain boundary phase. However, the effect of the present invention cannot be enjoyed only when the Zr-enriched product exists in the grain boundary phase.

Ti is conventionally known as an additive element to form a product in the R ₂ T ₁₄ B phase in RTB-based rare earth permanent magnets (for example, J. Appl. Phys. 69 (1991) 6055). The inventors of the present invention found that adding Zr and Ti to form a product in the R ₂ T ₁₄ B phase is effective in expanding the sintering temperature range. Here, when Zr is added, even if it is added in an amount that can sufficiently expand the effect of the sintering temperature range, the magnetic properties are hardly lowered, specifically, the residual magnetic flux density (Br) is hardly lowered. On the other hand, it is known that when Ti is added, the residual magnetic flux density (Br) decreases remarkably when added in an amount sufficient to increase the effect of the sintering temperature range, which is not preferable in practice. As described above, by determining the composition of the product to be a Zr-rich composition, it becomes possible to stably produce a high-performance permanent magnet over a wide temperature range.

The inventors of the present invention confirmed that there are several points in the production method in order to make the Zr-enriched product exist in the R ₂ T ₁₄ B phase. A series of steps in the method for producing a permanent magnet according to the present invention will be described later, but the necessary conditions for the Zr-enriched product to exist in the R ₂ T ₁₄ B phase will be described here.

As the manufacturing method of RTB series rare earth permanent magnets, there are two methods: the method of using a single alloy consistent with the required composition as the initial raw material (hereinafter referred to as "single method"), and the method of using multiple alloys with different compositions. The method in which the alloy is the starting material (hereinafter referred to as "mixing method"). The hybrid method typically uses an alloy mainly containing R ₂ T ₁₄ B phase (low R alloy) and an alloy containing more R than the low R alloy (high R alloy) as starting materials.

The inventors of the present invention obtained an RTB-based rare earth permanent magnet by adding Zr to either a low-R alloy or a high-R alloy. As a result, it was confirmed that when a permanent magnet was produced by adding Zr to the low-R alloy, a Zr-enriched product existed in the R ₂ T ₁₄ B phase. On the other hand, it was also confirmed that in the case of producing a permanent magnet by adding Zr to a high-R alloy, the Zr-enriched product does not exist in the R ₂ T ₁₄ B phase.

Also, even when Zr is contained in the low R alloy, if the Zr-enriched product exists in the R ₂ T ₁₄ B phase at the stage of the low R alloy, the Zr-enriched product exists in the sintered phase after sintering. The triple point in the structure is enriched in the R phase (grain boundary phase), but the existence of Zr-enriched products in the R ₂ T ₁₄ B phase has not been confirmed. Therefore, in order to have Zr-enriched products in the R ₂ T ₁₄ B phase of the RTB-based rare earth permanent magnet, it is important to prevent the Zr-enriched products in the R ₂ T ₁₄ B phase at the raw material alloy stage.

For this reason, it is necessary to consider the manufacturing method of the raw material alloy. When producing a low R alloy by strip casting, it is necessary to control the peripheral speed of the cooling roll. When the peripheral speed of the cooling roll is low, a Zr-enriched product is formed in the R ₂ T ₁₄ B phase of the low R alloy while causing the precipitation of α-Fe. According to the research of the present inventors, the peripheral speed of the cooling roll is in the range of 1.0 to 1.8 m/s, and a low R alloy without Zr-enriched products in the R ₂ T ₁₄ B phase can be obtained. Furthermore, a permanent magnet with high magnetic properties can be obtained by using this low-R alloy.

Also, even if a low-R alloy in which no Zr-enriched product exists in the R ₂ T ₁₄ B phase is obtained, it is not preferable for the present invention to use it as a raw material alloy after heat treatment. This is because a Zr-enriched product is formed in the R ₂ T ₁₄ B phase of the low R alloy by heat treatment in a temperature range (approximately 700° C. or higher) that changes the structure of the low R alloy.

<chemical composition>

Next, the ideal chemical composition of the R-T-B based rare earth permanent magnet of the present invention will be described. The chemical composition mentioned here refers to the chemical composition after sintering.

The rare earth permanent magnet of the present invention contains 25 to 35% by weight of R.

Here, R is one or more selected from La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Yb, Lu, and Y. When the R content is less than 25% by weight, the formation of the R ₂ T ₁₄ B ₁ phase which is the main phase of the rare earth permanent magnet is insufficient. Therefore _, α-Fe with soft magnetic properties is precipitated, and the _coercive force _is significantly reduced; Density decreases. Also, when the R content exceeds 35% by weight, R reacts with oxygen to increase the amount of oxygen contained, and accordingly the R-rich phase effective for generating coercive force decreases, resulting in a decrease in coercive force. Therefore, the amount of R is determined to be 25 to 35% by weight. A preferable amount of R is 28 to 33% by weight, and a more preferable amount of R is 29 to 32% by weight.

Nd is rich in resources and relatively cheap, so it is ideal to choose Nd as the main component of R. Also, Dy is effective for increasing the anisotropic magnetic field of the R ₂ T ₁₄ B ₁ phase and improving the coercive force. Therefore, Nd and Dy are selected as R, and the total of Nd and Dy is preferably 25 to 33% by weight. Furthermore, within this range, the amount of Dy is preferably 0.1 to 8% by weight. It is preferable to determine the amount of Dy within the above-mentioned range according to the respective degrees of emphasis on the remanence magnetic flux density and the coercive force. That is, when a high residual magnetic flux density is desired, the amount of Dy is preferably 0.1 to 3.5% by weight, and when a high coercive force is desired, the amount of Dy is preferably 3.5 to 8% by weight.

Also, the rare earth permanent magnet of the present invention contains 0.5 to 4.5% by weight of boron (B). When B is less than 0.5% by weight, a high coercive force cannot be obtained; however, when B exceeds 4.5% by weight, the residual magnetic flux density tends to decrease. Therefore, the upper limit is 4.5% by weight. A preferable B content is 0.5 to 1.5% by weight, and a more preferable B content is 0.8 to 1.2% by weight.

The R-T-B based rare earth permanent magnet of the present invention can contain one or both of Al and Cu in a range of 0.02 to 0.6% by weight. By containing one or both of Al and Cu within this range, it becomes possible to increase the coercive force, increase the corrosion resistance, and improve the temperature characteristics of the obtained permanent magnet. When adding Al, the preferable amount of Al is 0.03 to 0.3% by weight, and the more preferable amount of Al is 0.05 to 0.25% by weight. Also, when Cu is added, the amount of Cu is 0.3% by weight or less (excluding 0), preferably 0.15% by weight or less (excluding 0), and more preferably 0.03 to 0.08% by weight.

The RTB-based rare earth permanent magnet of the present invention preferably contains 0.03 to 0.25% by weight of Zr in order to form a Zr-enriched product in the R ₂ T ₁₄ B phase. In order to improve the magnetic properties of RTB-based rare earth permanent magnets, Zr exerts the effect of inhibiting the abnormal growth of grains during sintering when the oxygen content is reduced, and makes the structure of the sintered body uniform and fine. Therefore, Zr has a remarkable effect when the oxygen content is low. The preferable content of Zr is 0.05 to 0.2 weight%, and the more preferable content is 0.1 to 0.15 weight%.

The R-T-B series rare earth permanent magnet of the present invention has an oxygen content below 2000ppm. When the oxygen content is high, the oxide phase which is a non-magnetic component increases, and the magnetic properties are lowered. Here, in the present invention, the amount of oxygen contained in the sintered body is determined to be 2000 ppm or less, preferably 1500 ppm or less, more preferably 1000 ppm or less. However, simply reducing the oxygen content reduces the oxide phase which has the effect of suppressing grain growth, and tends to cause grain growth in the process of obtaining a sufficient increase in density during sintering. Here, in the present invention, the R-T-B system rare earth permanent magnet contains Zr in a predetermined amount, which exhibits the effect of suppressing abnormal grain growth during sintering.

The R-T-B series rare earth permanent magnet of the present invention contains Co below 4% by weight (excluding 0), preferably 0.1-2.0% by weight, more preferably 0.3-1.0% by weight. Co forms the same phase as Fe, and is effective in raising the Curie temperature and improving the corrosion resistance of the grain boundary phase.

<Manufacturing method>

Next, a suitable method for producing the R-T-B-based rare earth permanent magnet of the present invention will be described.

In this embodiment, the RTB-based rare-earth permanent magnet of the present invention is manufactured using an alloy (low R alloy) mainly containing R ₂ T ₁₄ B phase and an alloy containing more R than the low R alloy (high R alloy). method to describe.

First, low R alloys and high R alloys are obtained by carrying out strip continuous casting of raw metals in vacuum or inert gas, preferably in an Ar protective atmosphere. Here, as described above, it is necessary to consider that no Zr-enriched product is generated in the R ₂ T ₁₄ B phase for the strip to be obtained, especially the low-R alloy strip. Specifically, the peripheral speed of the cooling roll is set in the range of 1.0 to 1.8 m/s. The ideal peripheral speed of the cooling roll is 1.2 to 1.5 m/s.

After a low-R alloy having an R ₂ T ₁₄ B phase free of Zr-rich products is obtained, the products are not allowed to form in the R ₂ T ₁₄ B phase until the sintering process described later, even if the product is maintained. The morphology of the R ₂ T ₁₄ B phase is important for the present invention. For example, before the pulverization process starting from hydrogen pulverization, it is desirable to avoid the heat treatment of heating and holding the low-R alloy at 700° C. or higher. This point will be further described in the first embodiment described later.

In the form of this embodiment, the characteristic matter is that Zr is added to the low R alloy. As explained in the column of <Structure>, this is because adding Zr to a low-R alloy without Zr-enriched products in the R ₂ T ₁₄ B phase can make the R ₂ T ₁₄ of the RTB-based rare earth permanent magnet Zr-enriched products exist in the B phase. In addition to rare earth elements, Fe, Co, and B, Cu and Al can be contained in the low-R alloy. In addition, in addition to rare earth elements, Fe, Co, and B, Cu and Al can also be contained in the high R alloy.

After the low-R alloy and the high-R alloy are produced, their raw material alloys can be pulverized separately or together. The pulverization process includes a coarse pulverization process and a fine pulverization process. First, the raw material alloys are roughly pulverized to a particle diameter of several hundreds of micrometers. Coarse pulverization is preferably carried out in an inert protective gas with a pounder, jaw crusher, Brown mill, or the like. In order to improve the coarse pulverization property, it is effective to perform coarse pulverization after absorbing hydrogen. In addition, hydrogen may be released after hydrogen absorption, followed by coarse pulverization.

After the coarse pulverization process is completed, it moves to the fine pulverization process. Fine pulverization mainly uses a jet mill, and the coarse powder with a particle diameter of about several hundred μm is crushed to an average particle diameter of 3 to 5 μm. The jet mill sprays high-pressure inert gas (such as nitrogen) from a narrow nozzle to generate a high-speed airflow, and the high-speed airflow accelerates the coarsely pulverized powder to cause the coarsely pulverized powder to collide with each other, and to collide with the target or The method of crushing by the impact of the container wall.

In the fine pulverization step, when the low R alloy and the high R alloy are pulverized separately, the finely pulverized low R alloy powder and the high R alloy powder are mixed in a nitrogen atmosphere. The mixing ratio of the low-R alloy powder and the high-R alloy powder may be about 80:20-97:3 by weight. Similarly, when the low-R alloy powder and the high-R alloy powder are pulverized together, the mixing ratio may be about 80:20 to 97:3 by weight. At the time of fine pulverization, by adding additives such as zinc stearate in an amount of about 0.01 to 0.3% by weight, a fine powder with high orientation during molding can be obtained.

Next, the mixed powder composed of low R alloy powder and high R alloy powder is filled into a mold surrounded by electromagnets, and a magnetic field is applied to make the crystal axes in an oriented state and formed in the magnetic field. Molding in this magnetic field may be carried out in a magnetic field of 12.0 to 17.0 kOe at a pressure of about 0.7 to 1.5 t/cm ² .

After forming in a magnetic field, its formed body is sintered in vacuum or in an inert gas protective atmosphere. It is necessary to adjust the sintering temperature according to various conditions such as composition, crushing method, particle size and particle size distribution, and sintering at 1000-1100°C for 1-5 hours is sufficient. In the present invention, a Zr-enriched product is generated in the R ₂ T ₁₄ B phase in the sintering step. Although the mechanism of the formation of Zr-enriched products that do not exist in the low-R alloy stage after sintering is not clear, the Zr solid-dissolved in the R ₂ T ₁₄ B phase in the low-R alloy stage is precipitated during the sintering process. Possibilities exist.

After sintering, aging treatment may be applied to the obtained sintered body. Aging treatment is important in controlling the coercive force. When the aging treatment is performed in two stages, it is effective to carry out the heat preservation at around 800°C and around 600°C for a predetermined time. The coercivity increases when heat treatment near 800°C is performed after sintering, so the mixing method is particularly effective. In addition, since the coercive force greatly increases during heat treatment at around 600°C, when aging treatment is performed in one stage, it is only necessary to perform aging treatment at around 600°C.

(Example)

<First embodiment>

The R-T-B series rare earth permanent magnets were produced according to the following production process.

1) Raw material alloy

A raw material alloy (strip material) having the composition and thickness shown in FIG. 1 was produced by the strip continuous casting method. Peripheral speed of cooling roll: 1.5m/s for low R alloy; 0.6m/s for high R alloy. However, for the low R alloy of Comparative Example 3 in Fig. 1, the peripheral speed of the cooling roll was set at 0.6 m/s. The thickness of the alloy is the average value of the thicknesses of 50 cast pieces (strips). Also, for the low R alloy in Example 1 shown in FIG. 1 , no Zr-enriched product (hereinafter referred to as an in-phase product) was found in the R ₂ T ₁₄ B phase. In contrast, in Comparative Example 3 It was confirmed that Zr-enriched products existed in the R ₂ T ₁₄ B phase in the low R alloy.

2) Hydrogen crushing process

After absorbing hydrogen at room temperature, dehydrogenation was performed at 600° C. for 1 hour in an Ar protective atmosphere, and hydrogen pulverization treatment was performed.

In order to obtain high magnetic properties, the oxygen content of the sintered body was controlled below 2000ppm in this test, so the protective atmosphere of each process from hydrogen pulverization (recovery after pulverization) to sintering (putting into the sintering furnace) was controlled to less than 100ppm. oxygen concentration.

3) Mixing-crushing process

Generally, two-stage pulverization of coarse pulverization and fine pulverization is performed, but the coarse pulverization step is omitted in this example.

Add 0.05% by weight of zinc stearate before finely pulverizing, with the combination of embodiment 1 shown in Figure 1, comparative example 1～comparative example 3, low R alloy and high R alloy are mixed in a screw mixer (nota) mixer) for 30 minutes. In addition, in any of Example 1 and Comparative Examples 1 to 3, the mixing ratio of the low-R alloy and the high-R alloy was 90:10.

Then, it was finely pulverized with a jet mill until the average particle diameter was 4.8 to 5.1 μm.

4) Forming process

The obtained fine powder was molded in a magnetic field of 15.0 kOe at a pressure of 1.2 t/cm ² to obtain a molded body.

5) Sintering and aging process

This molded body was sintered at 1070° C. for 4 hours in a vacuum and then quenched. Next, the obtained sintered body was subjected to two-stage aging treatment at 800° C. for 1 hour and 550° C. for 2.5 hours (both in an Ar protective atmosphere).

The magnetic properties of the obtained permanent magnets were measured with a BH tracer, and the results are shown in FIGS. 2 to 5 . In addition, in FIGS. 2 to 5 , Br represents the residual magnetic flux density, HcJ represents the coercive force, and "Hk/HcJ" represents the squareness ratio. Also, the squareness ratio (Hk/HcJ) is an index of magnetic performance, and indicates the degree of squareness in the second quadrant of the hysteresis loop. Also, Hk is the external magnetic field intensity at which the magnetic flux density in the second quadrant of the hysteresis loop becomes 90% of the residual magnetic flux density. In FIGS. 2 to 5 , those in which the presence of in-phase products were confirmed are indicated by "◯", and those in which the absence of in-phase products were confirmed are indicated by "×". Confirmation of the in-phase product was based on observation with TEM (Transmission Electron Microscope: transmission electron microscope (TEM-3010 manufactured by JEOL Ltd.)). The observation sample was made by ion milling method, and the C surface of the R ₂ T ₁₄ B phase was observed. The chemical composition of the obtained sintered body is shown in the column of "Sintered body composition" in FIG. 1 . Also, in Comparative Example 3, no in-phase products were confirmed, but Zr-enriched products were confirmed in the grain boundary phase.

It can be seen from Fig. 2 and Fig. 5 that for the R-T-B series rare earth permanent magnets (Example 1 and Comparative Example 1) in which intraphase products were confirmed, abnormal grain growth was suppressed, and by adding a small amount of M (Zr or Ti), The squareness ratio (Hk/HcJ) is improved. However, as shown in FIG. 3 , when Ti is selected as the additive element M, the residual magnetic flux density (Br) decreases significantly. Also, for the R-T-B series rare earth permanent magnets (Comparative Example 2 and Comparative Example 3) that were confirmed to have no in-phase products, the squareness ratio (Hk/HcJ) was improved by adding a large amount of 0.2% by weight of Zr (see FIG. 5 ), But the residual magnetic flux density (Br) is still greatly reduced (refer to Figure 3). As described above, it has been confirmed that the R-T-B based rare earth permanent magnet having the in-phase product can obtain a high squareness ratio (Hk/HcJ) while suppressing a decrease in the residual magnetic flux density (Br).

In addition, for Comparative Example 3, in which it was confirmed that there were in-phase products in the R ₂ T ₁₄ B phase in the low R alloy stage, the reason why there were no in-phase products in the RTB-based rare earth permanent magnet was presumed as follows: The Zr-enriched product (in-phase product) generated in the R ₂ T ₁₄ B phase grows very coarsely. It is presumed that this product does not cause volume expansion even after hydrogen pulverization treatment, so it can be explained that cracks are generated at the interface between the R ₂ T ₁₄ B phase and this product during hydrogen pulverization. When subjected to the pulverization step in this state, the product is separated from the R ₂ T ₁₄ B phase, and as a result, the product is no longer included in the R ₂ T ₁₄ B phase and exists independently from the R ₂ T ₁₄ B phase. Therefore, it can be considered that the RTB-based rare earth permanent magnet according to Comparative Example 3 has Zr-enriched products only in the grain boundary phase even after the sintering process.

Regarding the RTB-based rare earth permanent magnet having a Zr content of 0.10% by weight according to Example 1, TEM observation was performed in the same manner as above. The observation results are shown in FIGS. 6 to 8 . Also, Fig. 6 is a TEM photo of a sample with a Zr content of 0.10% by weight, and Fig. 7 is an EDS (Energy Dispersive X-ray Fluorescence Spectroscopy: Energy _{Dispersive X} -ray Fluorescence Spectroscopy: Energy dispersive X-ray analysis device (spectrometry) distribution diagram; Figure 8 is a high-resolution TEM photo of the sample.

As shown in FIG. 6 , in the R ₂ T ₁₄ B phase, an in-phase product with a relatively large axis was confirmed. The product has a sheet-like (ie plate-like) or needle-like form. In addition, since FIG. 6 is a photograph of the cross section of the observed sample, it is difficult to determine whether the in-phase product is flake-like or needle-like. Considering the observation results of other samples and FIG. 8 , the in-phase product has a length of several 100 μm and a width of several nm to 15 nm. The detailed chemical composition of the product in this phase is unclear, but it can be confirmed from Fig. 7(a) that the product in this phase is at least enriched in Zr. Also, as a result of observation of other samples, in addition to the in-phase products with large axial ratios, amorphous and circular in-phase products were also observed as shown in FIGS. 9 and 10 . Also, in Example 1, as a result of observing 20 crystal grains (R ₂ T ₁₄ B phase), in-phase products were observed in 6 crystal grains. In contrast, in Comparative Example 2, no in-phase product was observed for all 20 crystal grains (R ₂ T ₁₄ B phase).

The lower part of Fig. 11(a) shows the Zr mapping results of EPMA (Electron Probe Micro Analyzer: Electron Probe Micro Analyzer) of the sample whose Zr amount is 0.10% by weight in Example 1. The upper part of FIG. 11( a ) shows a composition image in the same field of view as the Zr mapping result shown in the lower part of FIG. 11( a ). Also, the lower part of FIG. 11( b ) shows the results of Zr mapping by EPMA of the sample of Comparative Example 2 in which the amount of Zr is 0.10% by weight. The upper part of Fig. 11(b) shows a composition image in the same field of view as the Zr mapping result shown in the lower part of Fig. 11(b).

Similar to the TEM observation results, it can be seen from Fig. 11(a) that the Zr-enriched R ₂ T ₁₄ B phase exists in Example 1, and Zr also exists in the grain boundary phase. In contrast, from FIG. 11( b ), it was not confirmed that the Zr-enriched R ₂ T ₁₄ B phase existed in Comparative Example 2, and Zr existed only in the grain boundary phase.

<Second embodiment>

Samples with 0.10% by weight of the added element M (Zr or Ti) in the composition of the sintered body were sintered at a temperature range of 1010 to 1090° C. for 4 hours, and R-T-B series rare earth permanent magnets were obtained in the same manner as in Example 1. . The magnetic properties of the obtained R-T-B based rare earth permanent magnets were measured in the same manner as in Example 1, and the results are shown in FIG. 12 . 13 to 15 show changes in magnetic properties with the sintering temperature. Also, FIG. 16 shows the magnetic properties at each sintering temperature plotted by the squareness ratio (Hk/HcJ) against the residual magnetic flux density (Br).

As shown in FIGS. 12 to 16 , it can be seen that when Zr is added as an additional element M to obtain an intraphase product, high magnetic properties can be stably obtained over a wide range of sintering temperatures. Specifically, according to Example 2 of the present invention, a residual magnetic flux density (Br) of 13.9 kG or more, a coercive force (HcJ) of 13.0 kOe or more, and a 95% or more Squareness ratio (Hk/HcJ). When Ti is added as an additional element M, the residual magnetic flux density (Br) decreases (Comparative Example 4), and when there are no intraphase products, the squareness ratio (Hk/HcJ) is not good, and the sintering temperature range is also narrow ( Comparative example 5).

<Third embodiment>

Four types of low-R alloys and two types of high-R alloys having the compositions and thicknesses shown in FIG. 17 were produced by strip casting at a cooling roll peripheral speed of 0.6 to 1.8 m/s. Then, four kinds of R-T-B series rare earth permanent magnets were obtained according to the combinations shown in FIG. 17 . Also, in any of samples A to D, the mixing ratio of the low R alloy and the high R alloy was 90:10. The low-R alloy and the high-R alloy shown in FIG. 17 were subjected to hydrogen pulverization in the same manner as in the first embodiment. After the hydrogen pulverization treatment, 0.05% by weight of butyl oleate was added, and the low-R alloy and the high-R alloy were mixed in a spiral mixer for 30 minutes according to the combination shown in FIG. 17 . Then it was finely pulverized by a jet mill to an average particle diameter of 4.1 μm. The obtained fine powder was formed in a magnetic field under the same conditions as in the first embodiment, and then sintered at 1010-1090° C. for 4 hours. Next, two-stage aging treatment of 800° C.×1 hour and 550° C.×2.5 hours was performed. The composition, oxygen content, and nitrogen content of the obtained sintered body are shown in FIG. 17 , and the magnetic properties are shown in FIG. 18 .

As shown in Figure 18, for sample A at a sintering temperature range of 1030-1070°C, a residual magnetic flux density (Br) of 14.0 kG or more, a coercive force (HcJ) of 13.0 kOe or more, and a rectangular shape of 95% or more can be obtained. Ratio (Hk/HcJ).

Compared with sample A, for sample B and sample C with lower Nd content, the residual magnetic flux density (Br) of 14.0kG or more and the coercivity of 13.5kOe or more can be obtained in the sintering temperature range of 1030-1070°C. Force (HcJ) and squareness ratio (Hk/HcJ) above 95%.

Compared with sample A, for sample D with higher Dy content, the residual magnetic flux density (Br) of 13.5kG or more and the coercive force (HcJ) of 15.5kOe or more can be obtained in the sintering temperature range of 1030-1070°C And a squareness ratio (Hk/HcJ) of 95% or more.

Furthermore, as a result of TEM observation of the samples sintered at 1050°C, in-phase products were observed in all samples.

From the above results, in the presence of intraphase products, high magnetic properties can be stably obtained over a wide sintering temperature range of 40°C or higher.

<Fourth embodiment>

Two kinds of low-R alloys and two kinds of high-R alloys were produced by the strip continuous casting method, and two kinds of R-T-B series rare earth permanent magnets were obtained according to the combinations shown in Fig. 19 . Also, for sample E, the mixing ratio of low R alloy and high R alloy was 90:10; on the other hand, for sample F, the mixing ratio of low R alloy and high R alloy was 80:20. The low-R alloy and the high-R alloy shown in FIG. 19 were subjected to hydrogen pulverization in the same manner as in the first embodiment. After the hydrogen pulverization treatment, 0.05% by weight of butyl oleate was added, and the low-R alloy and the high-R alloy were mixed in a screw mixer for 30 minutes according to the combination shown in FIG. 19 . Then it was finely pulverized by a jet mill to an average particle diameter of 4.0 μm. The obtained powder was molded in a magnetic field under the same conditions as in Example 1, and then sintered for sample E at 1070° C. for 4 hours and for sample F at 1020° C. for 4 hours. Next, samples E and F were subjected to two-stage aging treatment at 800° C.×1 hour and 550° C.×2.5 hours, respectively. The composition, oxygen content, and nitrogen content of the obtained sintered body are shown in FIG. 19 , and the magnetic properties are shown in FIG. 20 . For the convenience of comparison, the magnetic properties of samples A to D produced in the third embodiment are also shown in FIG. 20 .

Like samples A to F, even when the constituent elements are changed, a residual magnetic flux density (Br) of 13.8 kG or more, a coercive force (HcJ) of 13.0 kOe or more, and a squareness ratio (Hk/J) of 95% or more can be obtained. HcJ).

As described in detail above, by making Zr-enriched products exist in the R ₂ T ₁₄ B phase constituting the main phase of the RTB-based rare earth permanent magnet, it is possible to minimize the decrease in magnetic properties and suppress the crystal grains. grew up. In addition, according to the present invention, since a sintering temperature range of 40° C. or higher can be ensured, even in the case of using a large sintering furnace that tends to fluctuate in heating temperature, RTB-based rare earth permanent magnets with stable high magnetic properties can be easily obtained. magnet.

Claims (5)

1. R-T-B based rare earth element permanent magnet, it is by containing by R ₂T ₁₄The principal phase of B phase composition and contain than above-mentioned principal phase and more to many sintered body of crystal boundary phase of R and constitute, wherein R is more than a kind or a kind of rare earth element, rare earth element is that the notion, the T that contain Y is to be the essential transiens metallic element more than a kind or a kind with Fe or with Fe and Co, B is a boron element, and at above-mentioned R ₂T ₁₄There is the product of enrichment Zr in mutually in B.

2. the R-T-B based rare earth element permanent magnet of putting down in writing according to claim 1 is characterized in that, described product is sheet or needle-like.

3. the R-T-B based rare earth element permanent magnet of putting down in writing according to claim 1 is characterized in that, the oxygen amount that contains in the described sintered body is below the 2000ppm.

4. the R-T-B based rare earth element permanent magnet of putting down in writing according to claim 1, it is characterized in that the consisting of of described sintered body: R:28～33 weight %, B:0.5～1.5 weight %, Al:0.03～0.3 weight %, Cu:0.3 weight % following but do not comprise 0, Zr:0.05～0.2 weight %, Co:4 weight % following but do not comprise 0 and remainder be essentially Fe.

5. the R-T-B based rare earth element permanent magnet of putting down in writing according to claim 4 is characterized in that, the content of Zr is 0.1～0.15 weight % in the described sintered body.

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