Patent Grant
9412505
The present invention relates to a NdFeB system sintered magnet produced by a grain boundary diffusion treatment.
NdFeB system sintered magnets were discovered by Sagawa (one of the present inventors) and other researchers in 1982. NdFeB system sintered magnets exhibit characteristics far better than those of conventional permanent magnets, and can be advantageously manufactured from raw materials such as Nd (a kind of rare-earth element), iron, and boron, which are relatively abundant and inexpensive. Hence, NdFeB system sintered magnets are used in a variety of products, such as driving motors for hybrid or electric cars, battery-assisted bicycle motors, industrial motors, voice coil motors used in hard disks and other apparatuses, high-grade speakers, headphones, and permanent magnetic resonance imaging systems. NdFeB system sintered magnets used for those purposes must have a high coercive force HcJ, a high maximum energy product (BH)max, and a high squareness ratio SQ. The squareness ratio SQ is defined as Hk/HcJ, where Hk is the absolute value of the magnetic field when the magnetization value corresponding to a zero magnetic field is decreased by 10% on the magnetization curve extending across the boundary of the first and second quadrants of a graph with the horizontal axis indicating the magnetic field and the vertical axis indicating the magnetization.
One method for enhancing the coercive force of a NdFeB system sintered magnet is a “single alloy method”, in which Dy and/or Tb (the “Dy and/or Tb” is hereinafter represented by “RH”) is added to a starting alloy when preparing the alloy. Another method is a “binary alloy blending technique”, in which a main phase alloy which does not contain RH and a grain boundary phase alloy to which RH is added are prepared as two kinds of starting alloy powder, which are subsequently mixed together and sintered. Still another method is a “grain boundary diffusion method”, which includes the steps of creating a NdFeB system sintered magnet as a base material, attaching RH to the surface of the base material by an appropriate process, (such as application or vapor deposition), and heating the magnet to diffuse RH from the surface of the base material into the inner region through the boundaries inside the base material (Patent Document 1).
The coercive force of a NdFeB sintered magnet can be enhanced by any of the aforementioned methods. However, it is known that the maximum energy product decreases if RH is present in the main-phase grains inside the sintered magnet. In the case of the single alloy method, since RH is mixed in the main-phase grains at the stage of the starting alloy powder, a sintered magnet created from that powder inevitably contains RH in its main-phase grains. Therefore, the sintered magnet created by the single alloy method has a relatively low maximum energy product while it has a high coercive force.
In the case of the binary alloy blending technique, the largest portion of RH will be held in the boundaries of the main-phase grains. Therefore, as compared to the single alloy method, the technique can suppress the decrease in the maximum energy product. Another advantage over the single alloy method is that the amount of use of the rare metal, i.e. RH, is reduced.
In the grain boundary diffusion method, RH attached to the surface of the base material is diffused into the inner region through the boundaries liquefied by heat in the base material. Therefore, the diffusion rate of RH in the boundaries is much higher than the rate at which RH is diffused from the boundaries into the main-phase grains, so that RH is promptly supplied into deeper regions of the base material. By contrast, the diffusion rate from the boundaries into the main-phase grains is low, since the main-phase grains remain in the solid state. This difference in the diffusion rate can be used to regulate the temperature and time of the heating process so as to realize an ideal state in which the RH content is high only in the vicinity of the surface of the main-phase grains (grain boundaries) in the base material while the content of the same is low inside the main-phase grains. Thus, it is possible to further minimize the decrease in the maximum energy product (BH)max than in the case of the binary alloy blending technique while enhancing the coercive force. Another advantage over the binary alloy blending technique is that the amount of the rare metal, i.e. RH, used is reduced.
There are two kinds of methods for producing NdFeB system sintered magnets: a “press-applied magnet-production method” and a “press-less magnet-production method.” In the press-applied magnet-production method, fine powder of a starting alloy (which is hereinafter called the “alloy powder”) is put in a mold, and a magnetic field is applied to the alloy powder while pressure is applied to the alloy powder with a pressing machine, whereby the creation of a compression-molded body and the orientation of the same body are simultaneously performed. Then, the compression-molded body is removed from the mold and sintered by heating. In the press-less magnet-production method, alloy powder which has been put in a predetermined filling container is oriented, and sintered as it is held in the filling container, without undergoing the compression molding.
The press-applied magnet-production method requires a large-size pressing machine to create a compression-molded body. Therefore, it is difficult to perform the process in a closed space. By contrast, in the press-less magnet-production process, which does not use a pressing machine, the processes from the filling through the sintering can be performed in a closed space.
In the grain boundary diffusion method, the condition of the grain boundary significantly affects the way the RH, which is attached to the surface of the base material by deposition, application or another process, is diffused into the base material, such as how easily RH will be diffused and how deep it can be diffused from the surface of the base material. One of the present inventors has discovered that a rare-earth rich phase (i.e. the phase containing rare-earth elements in higher proportions than the main-phase grains) in the grain boundary serves as the primary passage for the diffusion of RH in the grain boundary diffusion method, and that the rare-earth rich phase is preferred to continuously exist, without interruption, through the grain boundaries of the base material in order to diffuse RH to an adequate depth from the surface of the base material (Patent Document 2).
A later experiment conducted by the present inventors has revealed the following fact: In the production of a NdFeB system sintered magnet, an organic lubricant is added to the alloy powder in order to reduce the friction between the grains of the alloy powder and help the grains easily rotate in the orienting process, as well as for other purposes. The lubricant contains carbon. Although the carbon contents are mostly oxidized during the sintering process and released to the outside of the NdFeB system sintered magnet, a portion of the carbon atoms remains inside the magnet. Among the remaining carbon atoms, those which remain in the grain boundary are cohered together, forming a carbon rich phase (a phase whose carbon content is higher than the average of the entire NdFeB system sintered magnet) in the rare-earth rich phase. The carbon atoms existing in the grain boundaries are more likely to be gathered at a grain-boundary triple point (a portion of the grain boundary surrounded by three or more main-phase grains), where the distance between the main-phase grains is large and impurities can easily gather, than in a two-grain boundary portion (a portion of the grain boundary sandwiched between two main-phase grains), where the distance between the main-phase grains is small and impurities cannot easily enter. Therefore, the largest portion of the carbon rich phase is formed at the grain-boundary triple point.
As already noted, the rare-earth rich phase existing in the grain boundary serves as the primary passage for the diffusion of RH into the inner region of the NdFeB system sintered magnet. Conversely, the carbon rich phase formed in the rare-earth rich phase acts like a weir which blocks the diffusion passage of RH and impedes the diffusion of RH through the grain boundary. If the diffusion of RH through the grain boundary is impeded, the RH content in the vicinity of the surface of the NdFeB system sintered magnet increases, and a larger amount of RH permeates the main-phase grains in the region in the vicinity of the surface, lowering the maximum energy product in that region. In some cases, in order to remove such a region having the lowered maximum energy product, the surface region of the NdFeB system sintered magnet is scraped off after the grain boundary diffusion treatment. However, this is a waste of the valuable element, RH.
Furthermore, since RH cannot be diffused across the entire magnet, the coercive force and the squareness ratio cannot be sufficiently improved.
The problem to be solved by the present invention is to provide a NdFeB system sintered magnet which is produced by the grain boundary diffusion method and yet has a high coercive force and squareness ratio with only a small decrease in the maximum energy product.
A NdFeB system sintered magnet according to the present invention aimed at solving the aforementioned problem is a NdFeB system sintered magnet having a base material produced by orienting powder of a NdFeB system alloy and sintering the powder, with Dy and/or Tb (RH) attached to and diffused from a surface of the base material through the grain boundary inside the base material by a grain boundary diffusion treatment, wherein the difference Cgx-Cx between the RH content Cgx (wt %) in the grain boundary and the RH content Cx (wt %) in main-phase grains which are grains constituting the base material at the same depth within a range from the surface to which RH is attached to a depth of 3 mm is equal to or larger than 3 wt %.
As already explained, when a carbon rich phase is formed at a grain-boundary triple point, the amount of inflow of RH into the grain-boundary triple point exceeds the amount of outflow of RH from the grain-boundary triple point, so that the RH content in that grain-boundary triple point increases. Due to the decrease in the amount of outflow of RH, the RH content in a two-grain boundary portion located farther than the grain-boundary triple point from the attachment surface becomes lower than the RH content in a two-grain boundary portion located closer to the attachment surface than the grain-boundary triple point. Therefore, in a conventional NdFeB system sintered magnet, there is a large difference in the RH content in the vicinity of the grain-boundary triple point, and RH is prevented from diffusing into deeper regions. An experiment conducted by the present inventors has demonstrated that, in conventional NdFeB system sintered magnets, the difference between the RH content in the grain boundary at a depth of 3 mm from the attachment surface and the RH content in the main-phase grains is approximately 1 wt %.
By contrast, in the NdFeB system sintered magnet according to the present invention, the difference in the RH content between the grain boundary and the main-phase grains is equal to or larger than 3% at least within a range from the surface to which RH is attached to a depth of 3 mm. From this fact, it can be said that RH is mainly diffused through the grain boundary, with only a smaller amount of RH permeating the main-phase grains. Therefore, the NdFeB system sintered magnet according to the present invention can achieve a higher coercive force and squareness ratio than the conventional NdFeB system sintered magnets by a grain boundary diffusion treatment while suppressing the amount of decrease in the maximum energy product.
In the production of the NdFeB system sintered magnet according to the present invention, for example, the percentage of the total volume of a carbon rich phase in a rare-earth rich phase at the grain-boundary triple points in the base material to the total volume of the rare-earth rich phase should preferably be equal to or lower than 50%. By using such a base material, it is possible to prevent RH from being blocked by the carbon rich phase during the grain boundary diffusion treatment, and to reduce the amount of RH permeating into the main-phase grains.
In the NdFeB system sintered magnet according to the present invention, RH is not localized in the vicinity of the surface but is evenly diffused in the grain boundaries of the entire magnet. Therefore, the NdFeB system sintered magnet according to the present invention can achieve a higher coercive force and squareness ratio than the conventional NdFeB system sintered magnets by a grain boundary diffusion treatment while suppressing the amount of decrease in the maximum energy product.
One example of the NdFeB system sintered magnet according to the present invention and its production method is hereinafter described.
A method for producing a NdFeB system sintered magnet according to the present example and a method according to a comparative example are hereinafter described by means of the flowcharts of
As shown in
The processes of Steps A3 through A5 are performed as a press-less process. The entire processes from Steps A1 through A5 are performed in an oxygen-free atmosphere.
As shown in
The temperature-programmed orientation is a technique in which the alloy powder is heated in the orienting process so as to lower the coercive force of each individual grain of the alloy powder and thereby suppress the mutual repulsion of the grains after the orientation. By this technique, it is possible to improve the degree of orientation of the NdFeB system sintered magnet after the production.
A difference between the method of producing a NdFeB system sintered magnet according to the present example and the method according to the comparative example is hereinafter described with reference to the temperature history of the hydrogen pulverization process.
By contrast, the method for producing a NdFeB system sintered magnet according to the present example does not use the thermal dehydrogenation. Therefore, as shown in
Thus, with the method for producing a NdFeB system sintered magnet according to the present example, it is possible to simplify the production process as well as significantly reduce the production time.
For each of the alloys having the compositions shown in Table 1 as Composition Numbers 1-4, the method for producing a NdFeB system sintered magnet according to the present example and the method for producing a NdFeB system sintered magnet according to the comparative example were applied. The results were as shown in Table 2.
Each of the results shown in Table 2 were obtained under the condition that the grain size of the alloy powder after the fine pulverization was controlled to be 2.82 μm in terms of D50 measured by a laser diffraction method. A 100 AFG-type jet mill manufactured by Hosokawa Micron Corporation was used as the jet mill for the fine pulverization process. A magnetic characteristics measurement device manufactured by Nihon Denji Sokki co., ltd (product name: Pulse BH Curve Tracer PBH-1000) was used for the measurement of the magnetic characteristics.
In Table 2, the data of “Dehydrogenation: No” and “Temperature-Programmed Orientation: No” show the results of the method for producing a NdFeB system sintered magnet according to the present example, while the data of “Dehydrogenation: Yes” and “Temperature-Programmed Orientation: Yes” show the results of the method for producing a NdFeB system sintered magnet according to the comparative example.
As shown in Table 2, when the thermal dehydrogenation was not performed, the pulverization rate of the alloy in the fine pulverization process was higher than in the case where the thermal dehydrogenation was performed, regardless of which composition of the alloy was used. This is probably because, in the case where the thermal dehydrogenation is performed, the structure inside the alloy which has been embrittled due to the hydrogen occlusion recovers its toughness as a result of the thermal dehydrogenation, whereas, in the case where the thermal dehydrogenation is not performed, the structure remains embrittled. Thus, the production method according to the present example in which the thermal dehydrogenation is not performed has the effect of reducing the production time as compared to the conventional method in which the thermal dehydrogenation is performed.
Although no temperature-programmed orientation was performed, the production method according to the present example achieved high degrees of orientation Br/Js which exceeded 95% and were comparable to the levels achieved by the production method according to the comparative example in which the temperature-programmed orientation was performed. A detailed study by the present inventors has revealed the fact that the magnetic anisotropy of the grains of the alloy powder (i.e. the coercive force of each individual grain) becomes lower in the case where the thermal dehydrogenation is not performed. When the coercive force of the individual grains is low, each grain will be a multi-domain structure in which reverse magnetic domains are formed along with the weakening of the applied magnetic field after the alloy powder has been oriented. As a result, the magnetization of each grain decreases, which alleviates the deterioration in the degree of orientation due to the magnetic interaction among the neighboring grains, so that a high degree of orientation is achieved. In principle, this is the same as what occurs during the process of improving the degree of orientation of a NdFeB system sintered magnet after the production is improved through the temperature-programmed orientation.
In summary, in the method for producing a NdFeB system sintered magnet according to the present example, although the temperature-programmed orientation is not performed, a high degree of orientation can be achieved as in the case of the temperature-programmed orientation, so that the production process can be simplified and the production time can be reduced.
Each of the sintering temperatures shown in Table 2 is the temperature at which the density of a sintered body for a given combination of the composition and the production method will be closest to the theoretical density of the NdFeB system sintered magnet. As shown in Table 2, it has been found that the sintering temperature in the present example tends to be lower than in the comparative example. The decrease in the sintering temperature leads to a decrease in the energy consumption through the production of the NdFeB system sintered magnet, and therefore, to the saving of energy. Another favorable effect is the extension of the service life of the mold, which is also heated with the alloy powder.
It can also been understood from the results of Table 2 that the NdFeB system sintered magnets produced by the method according to the present example have higher coercive forces HcJ than the NdFeB system sintered magnets produced by the method according to the comparative example.
Subsequently, a measurement by Auger electron spectroscopy (AES) was conducted to examine the fine structure of the NdFeB system sintered magnets produced by the method according to the present example as well as that of the NdFeB system sintered magnets produced by the method according to the comparative example. The measurement device was an Auger microprobe manufactured by JEOL Ltd. (product name: JAMP-9500F).
A brief description of the principle of the Auger electron spectroscopy is as follows: In Auger electron spectroscopy, an electron beam is cast onto the surface of a target object, and the energy distribution of Auger electrons produced by the interactions between the electrons and the atoms irradiated with those electrons is determined. An Auger electron has an energy value specific to each element. Therefore, it is possible to identify the elements existing on the surface of the target object (more specifically, in the region from the surface to a depth of a few nanometers) by analyzing the energy distribution of the Auger electrons (qualitative analysis). It is also possible to quantify the amounts of elements from the ratios of their peak intensities (quantitative analysis).
The distribution of the elements in the depth direction of the target object can be determined by an ion-sputtering of the surface of the target object (e.g. by a sputtering process using Ar ions).
The actual method of analysis was as follows: To remove contaminations from the surface of a sample, the sputtering of the sample surface was performed for 2-3 minutes before the actual measurement, with the sample inclined at an angle for the Ar sputtering (30 degrees from the horizontal plane). Next, an Auger spectrum was acquired at a few points of Nd-rich phase in the grain-boundary triple point where C and O could be detected. Based on the spectrum, a detection threshold was determined (ROI setting). The spectrum-acquiring conditions were 20 kV in voltage, 2×10−8 A in electric current, and 55 degrees in angle (from the horizontal surface). Subsequently, the actual measurement was performed under the same conditions to acquire Auger images for Nd and C.
In the present analysis, Auger images of Nd and C (
After the aforementioned regions were extracted, the total area of the Nd-rich grain-boundary triple-point region 11 and that of the C-rich areas 12 located in the Nd-rich grain-boundary triple-point region 11 were calculated. The calculated areas were defined as the volumes of the respective regions, and the ratio C/Nd of the two regions was calculated. Such an image processing and calculation was performed for each of a plurality of visual fields.
The surface of each of the NdFeB system sintered magnets of the present and comparative examples produced from Composition Number 2 were divided into small areas of 24 μm×24 μm, and the distributions of Nd and C as well as the C/Nd ratio were analyzed for each small area.
In the case of the NdFeB system sintered magnet of the present example, the C/Nd ratio was equal to or lower than 20% in most of the small areas. Although the C/Nd ratio reached 50% in some of the small areas, none of the small areas had a C/Nd ratio over 50%. The C/Nd ratio over the entire area (the entire group of the small areas) was 26.5%.
In the case of the NdFeB system sintered magnet of the comparative example, the C/Nd ratio was as high as 90% or even higher in almost all the small areas. The C/Nd ratio over the entire area was 93.1%.
In the following description, a NdFeB system sintered magnet in which the volume ratio of the C-rich regions to the Nd-rich grain-boundary triple-point regions is equal to or lower than 50% is called the “NdFeB system sintered magnet of the present example.” Furthermore, a NdFeB system sintered magnet which does not have this characteristic is called the “NdFeB system sintered magnet of the comparative example.”
The carbon content of the NdFeB system sintered magnet takes approximately the same value for each production method. The carbon content of a NdFeB system sintered magnet corresponding to Composition Number 3 in Table 1, which was measured by using the CS-230 type carbon-sulfur analyzer manufactured by LECO Corporation, was approximately 1100 ppm for a magnet produced by the method according to the comparative example and approximately 800 ppm for a magnet produced by the method according to the present example. A grain-size distribution of each of the NdFeB system sintered magnets produced by the method according to the present example was also determined by taking micrographs of the magnet within a plurality of visual fields (
Tables 3 and 4 show the magnetic characteristics of the NdFeB system sintered magnets of the present example and those of the NdFeB system sintered magnets of the comparative example, as well as their magnetic characteristics of after they have been employed as base materials for the grain boundary diffusion method.
Present Examples 1-4 in Table 3 are NdFeB system sintered magnets 1 which were respectively produced from the alloys of Composition Numbers 1-4 by the method according to the present example, each magnet measuring 7 mm in length, 7 mm in width and 3 mm in thickness, with the direction of magnetization coinciding with the thickness direction. Comparative Examples 1-4 in Table 4 are NdFeB system sintered magnets which were respectively produced from the alloys of Composition Numbers 1-4 by the method according to the comparative example, with the same size as Present Examples 1-4. Each of these NdFeB system sintered magnets of Present Examples 1-4 and Comparative Examples was used as a base material for the grain boundary diffusion method, as will be described later.
In this table, Br is the residual magnetic flux density (the magnitude of the magnetization J or magnetic flux B at a magnetic field of H=0 on the magnetization curve (J-H curve) or demagnetization curve (B-H curve)), Js is the saturation magnetization (the maximum value of the magnetization J), HcB is the coercive force defined by the demagnetization curve, HcJ is the coercive force defined by the magnetization curve, (BH)max is the maximum energy product (the maximum value of the product of the magnetic flux density B and the magnetic field H on the demagnetization curve), Br/Js is the degree of orientation, and SQ is the squareness ratio. Larger values of these properties mean better magnetic characteristics.
As shown in Table 3, when the composition is the same, the NdFeB system sintered magnet of the present example has a higher coercive force HcJ than the NdFeB system sintered magnet of the comparative example. There is no significant difference in the degree of orientation Br/Js. However, as for the squareness ratio SQ, the NdFeB system sintered magnets of the present example has achieved extremely high values as compared to the NdFeB system sintered magnets of the comparative example.
Table 4 below shows the magnetic characteristics after the grain boundary diffusion treatment was performed using each of the NdFeB system sintered magnets shown in Table 3 as the base material and using Tb as RH.
The grain boundary diffusion (GBD) treatment was performed as follows:
A TbNiAl alloy powder composed of 92 wt % of Tb, 4.3 wt % of Ni and 3.7 wt % of Al was mixed with a silicon grease by a weight ratio of 80:20. Then, 0.07 g of silicon oil was added to 10 g of the aforementioned mixture to obtain a paste, and 10 mg of this paste was applied to each of the two magnetic pole faces (7 mm×7 mm in size) of the base material.
After the paste was applied, the rectangular base material which was placed on a molybdenum tray provided with a plurality of pointed supports. The rectangular base material, being held by the supports, was heated in a vacuum of 10−4 Pa. The heating temperature was 880 degrees Celsius, and the heating time was 10 hours. Subsequently, the base material was quenched to room temperature, after which it was heated at 500 degrees Celsius for two hours and then once more quenched to room temperature.
As shown in Table 4, the magnets obtained by performing a grain boundary diffusion treatment using the NdFeB system sintered magnets of the present example as the base material had much higher coercive forces HcJ than the sintered magnets of the comparative example obtained by performing a grain boundary diffusion treatment using the NdFeB system sintered magnets of the comparative example as the base material. Furthermore, in the case where the NdFeB system sintered magnets of the comparative example were used as the base material, the squareness ratio SQ significantly deteriorated through the grain boundary diffusion treatment, whereas, in the case where the NdFeB system sintered magnets of the present example were used as the base material, the squareness ratio SQ barely deteriorated; it rather became higher in some cases.
The amounts of decrease in the maximum energy product (BH)max through the grain boundary diffusion treatment for the base materials of Present Examples 1-4 were 1.49 MGOe, 1.83 MGOe, 0.23 MGOe and 0.77 MGOe, respectively, while the values for the base materials of Comparative Examples 1-4 were 2.22 MGOe, 1.44 MGOe, 0.68 MGOe and 1.54 MGOe, respectively.
A comparison of these values demonstrates that, in the case of the NdFeB system sintered magnet of Present Example 2, the decrease in the maximum energy product after the grain boundary diffusion treatment was larger than that of the NdFeB system sintered magnet of Comparative Example 2 produced from the same starting alloy. However, in any of the other cases, the NdFeB system sintered magnet of the present example showed a smaller decrease in the maximum energy product than the NdFeB system sintered magnet of the comparative example produced from the starting alloy of the same composition. Furthermore, the amount of decrease was nearly one half of that of the comparative example.
Thus, in many cases, the NdFeB system sintered magnet of the present example undergoes a smaller decrease in the maximum energy product (BH)max after the grain boundary diffusion treatment than the NdFeB system sintered magnet of the comparative example produced from the starting alloy of the same composition.
The present inventors also measured the Tb content distribution in the grain boundary of the NdFeB system sintered magnet after the grain boundary diffusion treatment (which is hereinafter called the “GBD-treated magnet”), and particularly the Tb content distribution at the grain-boundary triple points and the two-grain boundary portions, for both the present example and the comparative example.
A comparison of the WDS mapping images of the GBD-treated magnet of the present example shown in
For each grain-boundary triple point in the GBD-treated magnets of the present example and the comparative example, the difference between the highest value of the Tb content at that grain-boundary triple point and the lowest value of the Tb content in the two-grain boundary portion leading to that grain-boundary triple point was calculated, and a histogram showing the content difference for each grain-boundary triple point was created. The result was as shown in
By contrast, in the case of the GBD-treated magnet of the comparative example (the result of “With Dehydrogenation Process” in
The present inventors also conducted a measurement on the diffusion of Tb in the depth direction from the Tb-application surface of each of the GBD-treated magnets of the present example and the comparative example.
In this measurement, the following processes were performed: Initially, a base material corresponding to Composition Number 2 (a sintered body before the grain boundary diffusion treatment) was oxidized except for one magnetic pole face. Subsequently, Tb was applied to the non-oxidized magnetic pole face, and the grain boundary diffusion treatment was performed. The NdFeB system sintered magnet after the grain boundary diffusion treatment (GBD-treated magnet) was cut at a plane perpendicular to the magnetic pole faces. A linear analysis of the Tb content was performed with an EPMA along a straight line parallel to the depth direction on the cut surface. The linear analysis was performed from the Tb-application surface to the opposite end under the same measurement conditions as described previously. For each sample, data were acquired along five lines spaced at intervals that could be resolved by the device. The five sets of data were superposed on each other to create a graph showing the Tb content in the depth direction. The conversion of data into the Tb content was performed by the same method as used for obtaining the images of
In each of the graphs of
As shown in
The difference Cs-Cd3 in the Tb content Cgx in the grain boundary between on the Tb-application surface (a depth of 0 mm) and at a depth of 3 mm from the Tb-application surface was equal to or larger than 25 wt % in the NdFeB system sintered magnet of the comparative example, while the difference was equal to or smaller than 20 wt % in the NdFeB system sintered magnet of the present example. Furthermore, the difference Cs-Cd1 in the Tb content Cgx in the grain boundary between on the Tb-application surface and at a depth of 1 mm from the Tb-application surface was equal to or larger than 20 wt % in the NdFeB system sintered magnet of the comparative example, while the difference was equal to or smaller than 15 wt % in the NdFeB system sintered magnet of the present example.
The difference in the Tb content between the main-phase grains and the grain boundary at a depth of 3 mm (where the content difference is the smallest) was approximately 1 wt % in the NdFeB system sintered magnet of the comparative example, whereas the same difference was equal to or larger than 3 wt % in the NdFeB system sintered magnet of the present example.
The results described thus far demonstrate that, as compared to the GBD-treated magnet of the comparative example, the GBD-treated magnet of the present example has a larger amount of Tb (RH) diffused in the depth direction, with only a smaller amount of Tb permeating the main-phase grains in the vicinity of the Tb-application surface. The large difference between the curves Cgx and Cx in
Indeed, the content Cx of Tb in the main-phase grains on the Tb-application surface of the GBD-treated magnet of the present example having the aforementioned characteristics was approximately 7 wt %, while it was approximately 12 wt % in the case of the GBD-treated magnet of the comparative example. This result confirms that the GBD-treated magnet of the present example has a smaller amount of Tb permeating the main-phase grains in the vicinity of the Tb-application surface than the GBD-treated magnet of the comparative example.
Therefore, in the GBD-treated magnet of the present example, the amount of decrease in the maximum energy product is smaller than in the GBD-treated magnet of the comparative example. The fact that the GBD-treated magnet of the present example has a higher coercive force and squareness ratio than the GBD-treated magnet of the comparative example is also probably due to the even diffusion of Tb in the grain boundary.
The fact that Tb can be diffused from one Tb-application surface to a depth of 3 mm suggests that, if Tb is applied to two opposite faces of a magnet, Tb can be diffused to the center of a GBD-treated magnet whose thickness is as large as 6 mm.
In the GBD-treated magnet of the present example, the low percentage of the carbon-rich phase in the Nd-rich phase of the sintered body used as the base material allows RH to be efficiently diffused through the Nd-rich phase in the grain boundaries. An experiment conducted by the present inventors has demonstrated that, when RH is applied to two opposite faces of a magnet, RH can be diffused to the center of a sintered base material whose thickness is as large as 10 mm. Table 5 shows an increase in the coercive force from the level before the grain boundary diffusion of the GBD-treated magnets of the present example corresponding to the alloys of Composition Numbers 1 and 3 as well as the GBD-treated magnet of the comparative example corresponding to the alloy of Composition Number 2, each of which was produced with three thicknesses of 3 mm, 6 mm and 10 mm.
As can be seen in this table, there is no significant difference between the GBD-treated magnets of the present example and that of the comparative example in the case of the 3-mm thickness. As the magnets become thicker, the GBD-treated magnets of the present example come to exhibit its superiority in terms of the coercive force. For example, in the case of the GBD-treated magnets of the present example, the amounts of increase in the coercive force at a thickness of 6 mm were maintained at approximately the same levels as they were at a thickness of 3 mm, whereas the amount significantly decreased in the case of the GBD-treated magnets of the comparative example. A larger increase in the coercive force suggests that RH is diffused to the center of the magnet. These results demonstrate that the GBD-treated magnets produced by the method according to the present example are suitable as a base material for producing a thick magnet having high magnetic characteristics by a grain boundary diffusion treatment.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2011-286864 | Dec 2011 | JP | national |
| 2012-026720 | Feb 2012 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2012/083789 | 12/27/2012 | WO | 00 |
| Publishing Document | Publishing Date | Country | Kind |
|---|---|---|---|
| WO2013/100011 | 7/4/2013 | WO | A |
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| Number | Date | Country | |
|---|---|---|---|
| 20140062631 A1 | Mar 2014 | US |
