Patent Application
20160240292
The present invention relates to a rare earth based permanent magnet, especially a rare earth based permanent magnet with part of R in the R-T-B based sintered magnet being replaced with heavy rare earth element(s).
The R-T-B based sintered magnet (R represents rare earth element(s), T represents Fe or Fe with part of it replaced by Co, and B represents boron) with the tetragonal compound R2T14B being its main phase is known to have excellent magnetic properties and thus is a representative permanent magnet with high performances since it was invented in 1982 (Patent Document 1).
The R-T-B based sintered magnet with the rare earth element(s) R being composed of Nd, Pr, Dy, Tb and/or Ho has a large magnetic anisotropy field Ha and is preferably used as a permanent magnet material. Especially the Nd—Fe—B based permanent magnet with Nd being the rare earth element R is widely used in consumer, industries, transportation equipments and the like because it has a good balance among the saturation magnetization Is, the Curie temperature Tc and magnetic anisotropy field Ha. Also, it is better than other R-T-B based sintered magnets in which other rare earth element(s) is/are used from the view point of the resource amount and corrosion resistance.
The improvement of magnetic properties is required in the conventional R-T-B based magnet. Particularly, a lot of efforts have been taken to improve the residual magnetic flux density Br and the coercivity HcJ. As one of the employed methods, a method is proposed that element(s) having high magnetic anisotropy such as Dy or Tb is/are added to increase the coercivity.
However, from the viewpoints of resource saving and cost reduction, the amount of the added heavy rare earth element(s) is required to be kept to a minimum. As the method for adding the heavy rare earth element(s), for example, a technique involving grain boundary diffusion has been disclosed (Patent Document 2).
As another method for adding the heavy rare earth element(s), a technique has been disclosed in which the RH-T phase (RH represents the heavy rare earth element) is mixed with the RL-T-B phase (RL represents the light rare earth element) or alternatively the RH-T-B phase is mixed with the RL-T-B phase to manufacture the sintered compact (Patent Document 3).
Patent Document 1: JP-A-S59-46008
Patent Document 2: JP-A-4831074
Patent Document 3: JP-A-4645855
In recent years, the utilization of the rare earth based magnet covers several aspects, and better magnetic properties compared to the conventional rare earth based magnet are desired. Especially when the R-T-B based sintered magnet is used in a hybrid vehicle or the like, the magnet is exposed to a relatively high temperature. Thus, the inhibition of the demagnetization at high temperature caused by heat becomes quite important. In order to inhibit the demagnetization at high temperature, the coercivity at room temperature needs to be increased in the R-T-B based sintered magnet.
The present invention is completed in view of the conditions above. For the R-T-B based sintered magnet, the present invention aims to provide a permanent magnet having a higher coercivity compared to that in the prior art.
In order to solve the technical problem mentioned above and reach the aim, the rare earth based permanent magnet of the present invention is characterized as follows. The rare earth based permanent magnet consists of a sintered compact having an R-T-B based composition, wherein the R contains R1 and R2 as the necessity (R1 represents at least one rare earth element including Y and excluding Dy, Tb and Ho, and R2 represents at least one from the group consisting of Dy, Tb and Ho). The rare earth based permanent magnet has a main phase grain with a core-shell structure which contains a core part and a shell part coating the core part. When the atom concentrations of R1 and R2 in the core part are defined as αR1 and αR2 respectively and the atom concentrations of R1 and R2 in the shell part are defined as βR1 and βR2 respectively, the following conditions are met, i.e., αR1<βR1, αR2>R2, αR1<αR2 and βR2<βR1. Further, relative to all the main phase grains observed at a unit cross-section of the sintered compact, the ratio occupied by the main phase grains having the core-shell structure is 5% or more.
In the present invention, a unit cross-section in the cross-section of the sintered compact is a region of 50 μm×50 μm.
In the R2T14B grain (the main phase grain), the part having a concentration difference in the heavy rare earth element of 3 at % or more compared with the outer edge part and containing the center of the main phase grain is defined as the core part, and the part of the main phase grain other than the core part is defined as the shell part. The main phase grain having the core part and the shell part is referred to as a core-shell grain. The part with a depth of 0.5 μm from the surface of the main phase grain is defined as the outer edge part, and the shell part contains the outer edge part.
The present inventors have studied whether the R-T-B based sintered magnet has a structure which can exert the high coercivity effect provided by the heavy rare earth element to the largest extent. As a result, it has been found that a high coercivity can be provided when the R-T-B based sintered magnet contains main phase grains having the core-shell structure mentioned above. The reason is not clear but is presumed by the present inventors as follows. First of all, the high coercivity is thought to be brought by the pinning effect of the magnetic domain wall generated at the interface between the core part and the shell part. Although the core part and the shell part have the same R2Fe14B structure, more R in the core part is the heavy rare earth element(s) and more R in the shell part is the light rare earth element(s). Thus, the lattice constants are different between the core part and the shell part. As a result, deformations are generated at the interface between the core part and the shell part. The deformations become the pinning sites, exerting the inhibitory effect on the movement of the magnetic domain wall. Secondly, the high coercivity is thought to be brought by the increased anisotropy magnetic field generated by the addition of the heavy rare earth element(s). Thirdly, the high coercivity is thought to be brought by the fact that the RL2T14B main phase (RL represents the light rare earth element(s) including Y) having fewer lattice defects coats the RH2T14B main phase (RH represents the heavy rare earth element) having more lattice defects. If defects such as the lattice defect are present on the surface of the main phase grain, they will become the nucleation site for magnetization reversal, resulting in the decrease of the coercivity. Thus, if the defects exist in a large number, the coercivity will decrease accordingly. Since the heavy rare earth element is easy to widely diffuse in the grain boundary phase and the energy in the RH2T14B main phase is less stable than that in the RL2T14B main phase, lattice defects are likely to occur in the RH2T14B main phase. In this way, when the RL2T14B main phase having fewer lattice defects is used to coat the RH2T14B main phase, the decrease of the coercivity caused by the lattice defects can be inhibited. In addition, when the ratio occupied by the main phase grains having the core-shell structure is 5% or more, the coercivity can be substantially increased.
In a preferable embodiment of the present invention, R2 contained in the sintered compact accounts for 11 at % or less.
When the content of the heavy rare earth element is 11 at % or less in the R-T-B based sintered magnet of the present invention, the substantial decrease of the residual magnetic flux density can be prevented. The reason why the residual magnetic flux density is decreased with the addition of the heavy rare earth element(s) is considered to be the decrease of magnetization, wherein the decrease of magnetization is caused by the anti-parallel coupling of the magnetic moment of the heavy rare earth element(s) and the magnetic momen of Nd or Fe. The present invention has been finished in view of the findings above.
As described above, the R-T-B based sintered magnet according to the present invention has a higher coercivity than the conventional ones.
Hereinafter, the present invention will be described in detail based on embodiments. However, the present invention is not limited to the following embodiments and examples. In addition, the constituent elements in the embodiments and examples described below include those can be easily thought of by those skilled in the art, those substantially the same and those with so-called equivalent scopes. Further, the constituent elements disclosed in the embodiments and examples described below can be properly used in combination or alternatively be appropriately selected.
The R-T-B based sintered magnet of the present embodiment contains 11 to 18 at % of the rare earth element(s) (R). If the concentration of R is less than 11 at %, the generation of R2T14B phases (which constitute the main phase of the R-T-B based sintered magnet) will not be complete and α-Fe or the like which possesses soft magnetism will be precipitated. Thus, the coercivity significantly decreases. On the other hand, if the content of R is higher than 18 at %, the volume ratio occupied by the R2T14B main phase decreases and the residual magnetic flux density will decrease. In addition, R reacts with oxygen, and thus the content of oxygen will increase. With this, the R-rich phase which helps the generation of coercivity will be less, leading to the decrease of the coercivity.
In the present embodiment, the rare earth element(s) (R) contains R1 and R2, wherein R1 represents at least one rare earth element including Y and excluding Dy, Tb and Ho, and R2 represents at least one from the group consisting of Dy, Tb and Ho. Preferably, relative to the total content of the rare earth element(s) (TRE), the ratio of R1 to TRE is 30 to 92 weight % and the ratio of R2 to TRE is 8 to 70 weight %. Here, R may also contain some other component(s) from the impurity of the raw material or the impurity mixed during manufacturing.
The R-T-B based sintered magnet of the present embodiment contains 5 to 8 at % of boron (B). When less than 5 at % of B is contained, no high coercivity can be provided. On the other hand, if more than 8 at % of B is contained, the residual magnetic flux density tends to decrease. Thus, the upper limit of B is set at 8 at %.
The R-T-B based sintered magnet of the present invention contains 74 to 83 at % of the transition metal element T. In the present invention, T contains Fe as the essential element and may contain 4.0 at % or less of Co. Co forms the same phase as Fe while it contributes to the increase of the Curie temperature and the improvement of corrosion resistance of the grain boundary phase. In addition, the R-T-B based sintered magnet which can be used in the present invention may contain either Al or Cu or both in an amount of 0.01 to 1.2 at %. If either Al or Cu or both is contained in such a range, the obtained sintered magnet can have a high coercivity, good corrosion resistance and improved temperature properties.
The R-T-B based sintered magnet of the present embodiment may contain other element(s). For example, the element such as Zr, Ti, Bi, Sn, Ga, Nb, Ta, Si, V, Ag, Ge or the like can be properly contained. On the other hand, it is preferable that the content of the impurity element(s) such as oxygen, nitrogen, carbon and the like is declined to the minimum. Especially for oxygen which is harmful to the magnetic properties, its content is preferably set at 5000 ppm or less and more preferably set at 3000 ppm or less. It is because that if the content of oxygen is high, the non-magnetic phase of oxides of the rare earth element(s) will increase, resulting in the deterioration of magnetic properties.
In the R-T-B based sintered magnet of the present embodiment, in addition to the R2T14B main phase, there is a complex structure composed of the eutectic compositions such as the R-rich phase, the B-rich phase and the like which are referred to as the grain boundary phase. The size of the main phase grains is approximately 1 to 10 μm.
Hereinafter, the preferable example of the manufacturing method in the present invention will be described.
During the manufacture of the R-T-B based sintered magnet of the present embodiment, alloy raw materials are prepared to provide the R1-T-B based magnet and the R2-T-B based magnet with desired compositions, respectively. The alloy raw materials can be manufactured by a strip casting method or other well-known melting methods under vacuum or in an inert atmosphere preferably Ar atmosphere. In the strip casting method, the metal raw material is melted under the nonoxidizing atmosphere such as Ar atmosphere and the obtained molten metal is sprayed to the surface of a rotating roll. The molten metal quenched on the roll will be solidified into a thin plate or a sheet (a scale-like shape). The quenched and solidified alloy is then provided with a homogeneous structure having a grain size of 1 to 50 μm. In addition to the strip casting method, the alloy raw material can also be obtained by some melting methods such as the high frequency induction melting method. In addition, in order to prevent the segregation from happening after the melting process, the molten metal can be poured onto a water-cooled copper plate so as to be solidified. Further, the alloy obtained by the reduction-diffusion method can be used as the alloy raw material.
The obtained R1-T-B based alloy raw material and the R2-T-B based alloy raw material are mixed and then subjected to the pulverization step. The mixing ratio can be properly adjusted in accordance with the target composition to be obtained after mixing or the like. Preferably, the weight ratio occupied by the R1-T-B based alloy is 30 to 92% and that occupied by the R2-T-B based alloy is 8 to 70%. The pulverization step includes a coarse pulverization step and a fine pulverization step. First of all, the alloy raw material is coarsely pulverized to have a particle size of approximately several hundreds of μm. The coarse pulverization is preferably performed in an inert atmosphere by using a stamp mill, a jaw crusher, a Braun mill or the like. Before the coarse pulverization, it is effective to perform the pulverization by storing hydrogen into the alloy raw material and then releasing the hydrogen. The hydrogen releasing treatment is performed to reduce the hydrogen which may turn to be an impurity for the rare earth based sintered magnet. The heating and holding temperature for hydrogen storage is set at 200° C. or higher and preferably 350° C. or higher. The holding time varies depending on the relationship with the holding temperature, the thickness of the alloy raw material and the like. However, it lasts for at least 30 minutes or longer and preferably for 1 hour or longer. The hydrogen releasing treatment is performed under vacuum or in an Ar gas flow. In addition, the hydrogen storing treatment and the hydrogen releasing treatment are not necessary treatments. Alternatively, the hydrogen pulverization can be deemed as the coarse pulverization, and thus the mechanical coarse pulverization can be omitted.
After the coarse pulverization, the alloy is transferred to the fine pulverization step. In the fine pulverization, a jet mill is mainly used to turn the coarsely pulverized powder having a particle size of several hundreds of μm into a powder with an average particle size of 2.5 to 6 μm and preferably 3 to 5 μm. The jet mill performs the following pulverization process. The jet mill ejects an inert gas with a high pressure through a narrow nozzle to provide a high-speeded gas flow. The coarsely pulverized powder is accelerated by this high-speeded gas flow, causing a collision between the coarsely pulverized powders or a collision between the coarsely pulverized powders and a target or the wall of a container.
A wet pulverization can also be used in the fine pulverization. In the wet pulverization, a ball mill or a wet attritor or the like can be used to turn the coarsely pulverized powder having a particle size of several hundred of μm into a powder with an average particle size of 1.5 to 5 μm and preferably 2 to 4.5 μm. In the wet pulverization, an appropriate dispersion medium is selected and the pulverization is performed with the powder of the magnet not contacting with oxygen. In this respect, a finely pulverized powder can be obtained with a low concentration of oxygen.
In order to improve the lubricity of the powder and help the powder to orient more easily in the pressing step, about 0.01 to 0.3 wt % of fatty acids or the derivatives thereof or hydrocarbons can be added during the fine pulverization. These fatty acids or the derivatives thereof or hydrocarbons can be, for example, zinc stearate, calcium stearate, aluminium stearate, Stearamide, Oleamide, ethylene bisstearamide which are the stearic acid-based or oleic acid-based compounds; paraffin and naphthalene which are hydrocarbons; and the like.
The fine powders mentioned above are subjected to a pressing step in a magnetic field. The pressure during the pressing in the magnetic field can be set to be 0.3 to 3 ton/cm2, i.e., 30 to 300 MPa. The pressure can be constant from the beginning to the end, or can be increasing or decreasing gradually, or can be changing irregularly. The lower the pressure is, the better the orientation will be. However, if the pressure is much too low, problems will arise during the handling due to insufficient strength of the green compact. From this point, the pressure should be selected from the range mentioned above. The final relative density of the green compact obtained by pressing in the magnetic field is usually 40 to 60%.
The magnetic field to be applied can be set at approximately 10 to 20 kOe, i.e., 960 to 1600 kA/m. The applied magnetic field is not limited to the static magnetic field, and it also can be a pulsed magnetic field. In addition, the static magnetic field and the pulsed magnetic field can be used in combination.
Then, the green compact is sintered under vacuum or in an inert gas atmosphere. The sintering temperature should be adjusted depending on the conditions such as the composition, the pulverization method, the average particle size, the particle size distribution and the like. In the present invention, the green compact is sintered at 850 to 950° C. With such a sintering temperature, the light rare earth element(s) will diffuse readily while the heavy rare earth element(s) is hard to diffuse. Thus, only the light rare earth element(s) diffuse widely. Further, in the outer edge part of the R22T14B main phase (R2 represents at least one from the group consisting of Dy, Tb and Ho), the light rare earth element(s) concentrates, and thus the desired structure can be obtained. If the sintering temperature is 1000° C. or higher, both the light rare earth element(s) and the heavy rare earth element(s) will diffuse widely, and thus no desired structure will be provided. Further, if the temperature is lower than 850° C., the temperature will be not sufficient for diffusion and thus the desired structure will not be obtained
The time for the sintering step should be adjusted depending on the conditions such as the composition, the pulverization method, the average particle size, the particle size distribution and the like. It is set as 48 to 96 hours. As the sintering temperature is low, the time should be 48 hours or longer so as to provide a sufficient density for the sintered compact. On the other hand, if the time is longer than 96 hours, the main phase grains grow, leading to a substantial decrease of the coercivity. The main phase grains in the sintered compact are preferably 10 μm or smaller in size.
After sintered, the obtained sintered compact can be subjected to an aging treatment. This step is crucial for the control of the coercivity. When the aging treatment is performed in two-step, it will be effective to last for a required time at about 800° C. and then about 600° C. respectively. If a heat treatment is performed at around 800° C. after the sintering step, the coercivity will increase. Thus, it is especially effective in the mixing method. In addition, as a heat treatment at around 600° C. greatly elevates the coercivity, the aging treatment can be performed at approximately 600° C. when the aging treatment is to be performed in one-step.
Hereinafter, the present invention will be described in detail based on the examples and comparative examples. However, the present invention is not limited to the following examples.
In order to prepare the R1-T-B based alloy and the R2-T-B based alloy, the metals or alloy raw materials were mixed together to provide raw materials having the compositions listed in Table 1. Then, they were melted and then casted by the strip casting method to provide alloy sheets respectively. In Examples 1 to 3, Dy, Tb and Ho were used as R2, respectively.
The obtained two kinds of alloy sheets were mixed in a weight ratio of 92:8 and then subjected to the hydrogen pulverization so as to provide the coarsely pulverized powders. Oleamide was added as the lubricant in an amount of 0.1 wt % into the coarsely pulverized powders respectively. Then, a jet pulverizer (a jet mill) was used to perform the fine pulverization under a high pressure in a nitrogen atmosphere respectively so that the finely pulverized powders were obtained.
Thereafter, the finely pulverized powders were put into a press mold and then pressed in the magnetic field. In specific, the pressing step was performed in a magnetic field of 15 kOe under a pressure of 140 MPa. In this respect, green compacts of 20 mm×18 mm×13 mm were obtained. The direction of the magnetic field was perpendicular to the direction in which the powders were pressed. The obtained green compacts were sintered at 850° C. for 48 hours. Then, they were provided with an aging treatment for 1 hour at 600° C.
The obtained sintered compacts were measured for the residual magnetic flux density (Br) and the coercivity (HcJ) by using a BH tracer. The results were shown in Table 3.
The obtained sintered compacts were cut down in a direction parallel to axis of easy magnetization and then resin-embedded into the epoxy resin. The cross-sections were polished using commercially available sandpapers, wherein the grit size of the sandpaper gradually became larger. At last, the cross-sections were polished by buff and diamond wheels. Here, the polishing step was performed without any water added. If water was used, the components in the grain boundary phase would be eroded.
The cross-sections of the sintered compacts were subjected to an ion milling to eliminate the influence of the oxide film or the nitride film on the outmost surface. Then, the cross-sections of the R-T-B based sintered magnet were observed by the EPMA (Electron Probe Micro Analyzer) and then analyzed. An area of 50 μm×50 μm was used as a unit cross-section and was subjected to the element mapping by EPMA (256 points×256 points). Here, the site to be observed in the cross-section was random. In this way, the main phase grains and the grain boundaries were determined. Also, to all of the main phase grains that can be identified in the unit cross-section area, it was determined that whether the core-shell structure was present. Further, the compositions of each core part and each shell part were determined.
The details for the method of analyzing the main phase grains were described as follows.
(1) According to the backscattered electron image obtained at the unit cross-section, the main phase grain part and the grain boundary part were identified by image analysis method.
(2) Based on the mapping data of the intensities of the characteristic x-ray of R1 and R2 obtained by EPMA, the element concentrations were calculated. The region containing the center of the main phase grain and having a concentration difference in the heavy rare earth element of 3% or more compared with the outer edge part of the main phase grain was defined as the core part, and the part other than the core part was defined as the shell part. For one visual field, the total grain number (D) and the number of the core-shell grains (E) were investigated. Then, the number ratio occupied by the core-shell grains (E/D) were calculated.
(3) The foregoing operations (1) and (2) were done in 20 visual fields in one cross-section of a single sample. In this way, the average concentration of the rare earth element in the core part of the core-shell grain (αR1 and αR2) and the average concentration of the rare earth element in the shell part of the core-shell grain (βR1 and βR2) were calculated. Then, the average value of the ratio occupied by the number of the core-shell grains per visual field was determined.
In order to prepare the R1-T-B based alloy, the metals or alloy raw materials were mixed together to provide the raw material having the composition as shown in Table 2. Then, they were melted and then casted by the strip casting method to provide alloy sheets.
The obtained alloy sheets were subjected to the hydrogen pulverization so as to provide a coarsely pulverized powder. Oleamide was added as the lubricant in an amount of 0.1 wt % into the coarsely pulverized powder. Then, a jet pulverizer (a jet mill) was used to perform the fine pulverization under a high pressure in a nitrogen atmosphere so that the finely pulverized powder was obtained.
Thereafter, the prepared R1-T-B based alloy powder was put into a press mold and then pressed in the magnetic field. In specific, the pressing step was performed in a magnetic field of 15 kOe under a pressure of 140 MPa. In this respect, a green compact of 20 mm×18 mm×13 mm was obtained. The direction of the magnetic field was perpendicular to the direction in which the powder was pressed. The obtained green compact was sintered at 1050° C. for 12 hours. Then, it was subjected to an aging treatment for 1 hour at 600° C. to provide a sintered compact.
The obtained sintered compact was measured similarly in Example 1 for Br and HcJ by using a BH tracer. The results were shown in Table 3.
In Examples 1 to 3, the main phase grains having a core-shell structure were obtained, wherein the core part had a higher atom concentration of the heavy rare earth element R2 and the shell part had a higher atom concentration of the light rare earth element R1. In addition, the coercivity was higher than that of Nd—Fe—B based sintered magnet from Comparative Example 1 where no heavy rare earth element was added. As described above, it was considered that the high coercivity was brought by the effects caused by the addition of the heavy rare earth element(s) and the presence of the core-shell structure, i.e., the increase of the magnetic anisotropy field, the pinning effect caused by the deformation and the reduction of the lattice defect-caused influence.
The preparation of the alloy sheets, pulverization, pressing, sintering and evaluation were similarly performed as in Example 1 except that Pr, Y, Ce or La was used as the light rare earth element(s) R1. The detailed compositions were listed in Table 4 and the evaluation results were shown in Table 5.
In Examples 4 to 7, the main phase grains having a core part and a shell part were present, wherein the core part had a higher amount of the heavy rare earth element(s) and the shell part had a higher amount of the light rare earth element(s). In addition, a high coercivity was provided. It can be determined that a core-shell structure and a high coercivity were similarly obtained as in Example 1 even if light rare earth element(s) other than Nd was/were introduced as R1.
In order to prepare the R1-T-B based alloy and the R2-T based alloy, the metals or alloy raw materials were mixed together to provide the raw materials having the composition listed in Table 6. Then, they were melted and then casted by the strip casting method to provide alloy sheets. The R1-T-B based alloy and the R2-T based alloy were mixed in a weight ratio of 93:7, and the pulverization, pressing, sintering and evaluation were similarly performed as in Example 1.
In order to prepare the R1-R2-T-B based alloy, the metals or alloy raw materials were mixed together to provide the raw materials having the composition listed in Table 6. Then, they were melted and then casted by the strip casting method to provide alloy sheets. The pulverization, pressing, sintering and evaluation were similarly performed as in Example 1.
In Comparative Example 2, a core-shell structure was formed, wherein the core part had a higher amount of the light rare earth element and the shell part had a higher amount of the heavy rare earth element. However, the coercivity was lower than that in Example 1. In Comparative Example 3, no core-shell structure could be found, and the coercivity was lower than that in Example 1.
The manufacture of the alloy sheets, pulverization, pressing, sintering and evaluation were similarly performed as in Example 1 except that the sintering temperature was different. In Comparative Examples 4 and 5. Examples 8 and 9, and Comparative Examples 6 and 7, the sintering temperatures were 750° C., 800° C., 900° C., 950° C., 1000° C. and 1050° C., respectively. The results were shown in Table 8.
In Example 1 and Examples 8 to 9, when the sintering temperature was set at 850 to 950° C., main phase grains were obtained in which the shell part had a higher amount of the light rare earth element R1 and the core part had a higher amount of the heavy rare earth element R2. The coercivity was higher than those in Comparative Examples 4 and 5 where no core-shell structure was formed. In Comparative Examples 4 and 5, no core-shell structure had been formed. This was probably due to that the sintering temperature is so low that the element R1 could hardly diffuse. In Comparative Examples 6 to 7, no core-shell structure can be formed at a temperature higher than 950° C., and thus the coercivity became lower. The reason for that might be that the sintering temperature was high and R1 widely diffused in the whole sintered compact.
The manufacture of the alloy sheets, pulverization, pressing, sintering and evaluation were similarly performed as in Example 1 except that the sintering time was different. In Comparative Examples 8 and 9, Examples 10 to 12, and Comparative Examples 10 and 11, the sintering time were 24 hours, 36 hours, 72 hours, 84 hours, 96 hours, 108 hours and 120 hours, respectively. The results were shown in Table 9.
In Examples 10 to 12, the same core-shell structure was present as that in Example 1, and a higher coercivity was provided. In Comparative Examples 8 and 9, no core-shell structure was present. As a result, the coercivity was low. The reason for this might be that the sintering time was short, and thus R had not sufficiently diffused. Besides, the residual magnetic flux density was also low. This was probably due to the fact that no sufficient sintering density could be provided, which is caused by that not only the sintering temperature was low but also the sintering time was short.
In Comparative Examples 10 and 11, the core-shell structure was similarly formed as in Example 1 but the coercivity was low. The decrease of the coercivity was probably due to the growth of the main phase grains resulting from the long sintering time.
The R1-T-B based alloys and the R2-T-B based alloys were similarly prepared as in Example 1. Then, they were mixed in an weight ratio of 98:2, 95:5, 92:8, 70:30, 50:50, 30:70, 20:80 and 10:90. And the pressing, sintering and evaluation were similarly performed as in Example 1. The compositions after mixing were shown in Table 10.
Then, Br and HcJ were similarly measured as in Example 1 by using a BH tracer. Thereafter, element mapping was performed by the EPMA. Also, the total number of main phase grains and the number of the core-shell grains were determined, and the concentrations of the rare earth element(s) in the core part and the shell part, i.e., αR1, αR2, βR1 and βR2, were determined. The results were shown in Table 11.
In all of Comparative Examples 12 to 13 and Examples 13 to 18, main phase grains with the structure formed by a core part and a shell part were contained. In particular, the core part had a higher amount of the heavy rare earth element(s), and the shell part had a higher amount of the light rare earth element(s). In addition, according to Examples 13 to 16, when the ratio occupied by the number of the core-shell grains was 5% or more and the concentration of R2 was 11 at % or less, the residual magnetic flux density was maintained to be high and a high coercivity was provided. In Comparative Examples 12 to 13 with less than 5% of the core-shell grains in number, the coercivity was low. It was considered that since a relatively low amount of the heavy rare earth element(s) was added, the number of the core-shell grains was small. Thus, the improving effect on the coercivity was not sufficient. In Examples 17 to 18 having the R2 concentration higher than 11 at %, a high coercivity was provided, but the residual magnetic flux density decreased greatly. This might be due to the addition of the heavy rare earth element(s), leading to the decreased saturation magnetization.
In order to prepare the R1-T-B based alloy and the R1-R2-T-B based alloy, the metals and the alloy raw materials were mixed together to provide the raw materials having the compositions as shown in Table 12. And they were melted and then casted by the strip casting method to provide alloy sheets respectively. Then, the pulverization, pressing, sintering and evaluation were similarly performed as in Example 1. The results were shown in Table 13.
In Examples 19 and 20, a core-shell structure was formed, wherein the core part had a higher amount of the heavy rare earth element(s) and the shell part had a higher amount of the light rare earth element(s). Further, compared to Comparative Example 1, a higher coercivity was provided. When compared to Example 1, a higher coercivity can be provided even if the ratio of R1 to R2 in the core part changed.
As described above, the R-T-B based sintered magnet of the present invention preserved a high residual magnetic flux density and a high coercivity. Thus, this magnet can be suitably used as the permanent magnet utilized in consumer, industries, transportation equipments and the like which require high output or high efficiency.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2015-027367 | Feb 2015 | JP | national |
| 2015-250286 | Dec 2015 | JP | national |
