Patent Grant
7550047
The present invention relates to a rare-earth-iron-boron based alloy, a magnetically anisotropic permanent magnet powder and a method of making such a powder, and an anisotropic bonded magnet including the magnetically anisotropic permanent magnet powder and a method of producing such a magnet.
A rare-earth-iron-boron based rare-earth magnet is a typical high-performance permanent magnet, has a structure including, as a main phase, an R2Fe14B-type crystalline phase, which is a ternary tetragonal compound, and exhibits excellent magnet performance. In R2Fe14B, R is at least one element selected from the group consisting of the rare-earth elements and yttrium and portions of Fe and B may be replaced with other elements.
Such rare-earth-iron-boron based rare-earth magnets are roughly classifiable into sintered magnets and bonded magnets. A sintered magnet is produced by compacting a fine powder of a rare-earth-iron-boron based magnet alloy (with a mean particle size of several μm) with a press machine and then sintering the resultant compact. On the other hand, a bonded magnet is usually produced by compacting a mixture (i.e., a compound) of a powder of a rare-earth-iron-boron based magnet alloy (with particle sizes of about 100 μm) and a binder resin within a press machine.
The sintered magnet is made of a powder with relatively small particle sizes, and therefore, the respective powder particles thereof exhibit magnetic anisotropy. For that reason, an aligning magnetic field is applied to the powder being compacted by the press machine, thereby obtaining a compact in which the powder particles are aligned with the direction of the magnetic field.
In the bonded magnet on the other hand, the powder particles used have particle sizes exceeding the crystal grain size, and normally exhibit no magnetic anisotropy and cannot be aligned under a magnetic field applied. Accordingly, to produce an anisotropic bonded magnet in which the powder particles are aligned with particular directions, a technique of making a magnetic powder, of which the respective powder particles exhibit the magnetic anisotropy, needs to be established.
To make a rare-earth alloy powder for an anisotropic bonded magnet, an HDDR (hydrogenation-disproportionation-desorption-recombination) process is currently researched and developed. The “HDDR” means a process in which hydrogenation, disproportionation, desorption and recombination are carried out in this order. In this HDDR process, a cast flake or powder of a rare-earth-iron-boron based alloy is maintained at a temperature of 500° C. to 1,000° C. within an H2 gas atmosphere or a mixture of an H2 gas and an inert gas so as to absorb hydrogen. As a result of this hydrogen absorption, the R2Fe14B phase is decomposed into rare-earth hydrides and iron-based borides. This reaction is represented by any of the following chemical equations:
R2Fe14B+2H22RH2+Fe2B+12Fe or
R2Fe14B+2H22RH2+Fe3B+11Fe
Thereafter, the hydrogenated flake or powder is subjected to a desorption process at a temperature of 500° C. to 1,000° C. and then cooled, thereby obtaining an alloy magnet powder. As a result of this desorption process, the R2Fe14B phase is regenerated from the hydrides or iron-based borides described above.
The respective R2Fe14B crystal grains, which had relatively large grain sizes (of several tens of μm or more, for example) before subjected to the hydrogenation process, turn into an aggregation of a huge number of very small R2Fe14B crystal grains (with grain sizes of approximately 0.1 μm to 1 μm). An aggregation of very small R2Fe14B crystal grains obtained in this manner will be referred to herein as a “recrystallized texture”. The very small R2Fe14B crystal grains in the recrystallized texture retain the crystallographic orientations of the original big R2Fe14B crystal grains. Accordingly, if the HDDR processed alloy powder is subjected to pulverization, classification and other processes such that its particle sizes are decreased to the sizes of the crystal grains yet to be subjected to the HDDR process or less, then the crystallographic orientations of those very small R2Fe14B crystal grains included in the respective powder particles can be aligned with a particular direction, thus realizing magnetic anisotropy. Also, the very small R2Fe14B crystal grains in the “recrystallized texture” have sizes that are close to the single domain critical grain size, thus achieving high coercivity, too.
Hereinafter, the HDDR process will be described with reference to
a) schematically illustrates a portion of a rare-earth-iron-boron based master alloy 1. Since the master alloy 1 is polycrystalline, there are a lot of grain boundaries 3 and not all of the crystallographic orientations 2 of its crystal grains are aligned with each other. Thus, the master alloy 1 is subjected to a coarse pulverization process, thereby forming powder particles 5, each of which is big enough to have a single crystallographic orientation as shown in
Next, the coarsely pulverized powder 4 is subjected to an HDDR process, thereby giving each particle 5 the recrystallized texture.
Subsequently, as shown in
A method of making a rare-earth-iron-boron based alloy powder with the recrystallized texture by performing such an HDDR process is disclosed in Japanese Patent Gazettes for Opposition No. 6-82575 and No. 7-68561, for example.
The magnetic powder prepared by the HDDR powder (which will be referred to herein as an “HDDR powder”), however, has the following drawbacks.
Firstly, to increase the remanence of the HDDR powder, the master alloy must be subjected to a homogenizing process to be carried out at an elevated temperature for a long time (e.g., at 1,100° C. for 20 hours). This process is required because if the master alloy has a fine texture, then the material powder yet to be subjected to the HDDR process will become polycrystalline and the powder particles will become magnetically isotropic.
Also, to subject the master alloy to the HDDR process in its entirety, hydrogen needs to be sufficiently diffused so as to reach the inside of the master alloy. To do so, the hydrogenation process must be carried out for a rather long time (e.g., at 800° C. for six hours). However, the longer the hydrogenation process time, the lower the saturation magnetization tends to be. The reason is as follows. Specifically, as the hydrogenation process lingers, the reversible reactions as represented by the above chemical formulae repeatedly occur an increasing number of times. Then, the crystallographic orientations of the R2Fe14B phase, which are retained in the master alloy, are gradually lost. As a result, the resultant “recrystallized texture” will have decreased magnetic anisotropy.
Nevertheless, if the hydrogenation process time is shortened, then the HDDR process will be incomplete, an insufficient amount of fine R2Fe14B phase will be produced, and both the coercivity HcJ and remanence Br will drop.
To overcome this problem, according to a proposed technique, Ga or any other element is added to the master alloy. In particular, by adding Ga to the master alloy, even if the hydrogenation process is carried out for a rather long time, the crystallographic orientations of the R2Fe14B phase, retained in the master alloy, are not lost so easily. As a result, both the coercivity HcJ and remanence Br can be increased to sufficient levels.
However, Ga is an expensive material, and the extended heat treatment for the purpose of hydrogenation increases the manufacturing cost, too. Thus, to mass-produce an HDDR powder at a reduced cost with the properties of the HDDR powder improved, the addition of that expensive Ga needs to be avoided and yet required magnet performance needs to be achieved in a short hydrogenation process time.
In Japanese Patent Gazettes for Opposition No. 6-82575 and No. 7-68561 identified above, a so-called alloy ingot, obtained by melting and casting a material alloy using an induction melting crucible, is used as the master alloy. Recently, however, a method for making a bonded magnet from a powder obtained by subjecting a strip-cast thin-plate material (or alloy flakes) to the HDDR process was also proposed (see Japanese Patent No. 3213638, for example).
However, the strip-cast alloy flakes include substantially no α-Fe phases, have a homogenous texture, but have excessively small crystal grain sizes. Thus, at a powder particle size applicable to a bonded magnet, the respective powder particles exhibit too low magnetic anisotropy to use the powder effectively.
In order to overcome the problems described above, a primary object of the present invention is to provide a rare-earth-iron-boron based alloy, which can eliminate the homogenizing process of the master alloy, can shorten the hydrogenation process time, and can improve both the coercivity HcJ and remanence Jr alike even substantially without adding Ga, and also provide a magnetically anisotropic magnet powder and a method of making the powder and an anisotropic bonded magnet and a method for producing the magnet.
A method of making a magnetically anisotropic magnet powder according to the present invention includes the steps of preparing a master alloy by cooling a rare-earth-iron-boron based molten alloy and subjecting the master alloy to an HDDR process. The step of preparing the master alloy includes the step of forming a solidified alloy layer, including a plurality of R2Fe14B-type crystals (where R is at least one element selected from the group consisting of the rare-earth elements and yttrium) in which rare-earth-rich phases are dispersed, by cooling the molten alloy through contact with a cooling member.
In one preferred embodiment, the step of forming the solidified alloy layer includes forming a first texture layer in contact with the cooling member and then further feeding the molten alloy onto the first texture layer to grow the R2Fe14B-type crystals on the first texture layer, thereby forming a second texture layer thereon.
In another preferred embodiment, the first texture layer consists essentially of R2Fe14B-type crystals with an average minor-axis size of less than 20 μm.
In another preferred embodiment, the R2Fe14B-type crystals of the second texture layer have an average minor-axis size of at least 20 μm and an average major-axis size of at least 100 μm.
As used herein, each region consisting of the R2Fe14B-type crystals in the alloy texture refers to a region with the same crystallographic orientation. The “region with the same crystallographic orientation” refers to herein a region showing the same contrast in an image obtained by observing a cross-sectional structure of the alloy with a polarizing microscope.
In another preferred embodiment, the solidified alloy layer includes the first and second texture layers, and the first texture layer accounts for less than 10 vol % of the overall solidified alloy layer.
In another preferred embodiment, the rare-earth-rich phase in the second texture layer is dispersed at an average interval of 50 μm or less in the R2Fe14B-type crystals.
In another preferred embodiment, the master alloy includes at most 5 vol % of α-Fe phase.
In another preferred embodiment, the rare-earth element included in the master alloy has a concentration of 26 mass % to 32 mass %.
In another preferred embodiment, Ga included in the master alloy has a concentration of 0.6 mass % or less.
In another preferred embodiment, the molten alloy is cooled at a rate of 10° C./s to 1,000° C./s and at a supercooling temperature of 100° C. to 300° C. in forming the first texture layer. In forming the second texture layer, the molten alloy is cooled at a rate of 1° C./s to 500° C./s.
In another preferred embodiment, the method includes the step of creating gaps in portions of the first texture layer that contact with the cooling member.
In another preferred embodiment, the molten alloy has a temperature of approximately 1,300° C. or less when reaching the cooling member.
In another preferred embodiment, the method includes the step of forming the solidified alloy layer by a centrifugal casting process.
In another preferred embodiment, the step of subjecting the master alloy to the HDDR process includes the step of heating the master alloy up to a temperature of 550° C. to 900° C. and then allowing the master alloy to react to hydrogen.
A rare-earth-iron-boron based alloy according to the present invention includes a first texture layer and a second texture layer, which is obtained by providing a plurality of R2Fe 14B-type crystals (where R is at least one element selected from the group consisting of the rare-earth elements and yttrium) where rare-earth-rich phases are dispersed, on the first texture layer. The first texture layer accounts for less than 10 vol% of the overall alloy. The R2Fe14B-type crystals have an average minor-axis size of 20 μm to 110 μm. The rare-earth-rich phases are dispersed at an average interval of 50 μm or less in the R2Fe14B-type crystals.
In one preferred embodiment, the alloy includes at most 5 vol % of α-Fe phase.
In another preferred embodiment, the rare-earth element has a concentration of 26 mass % to 32 mass %.
In another preferred embodiment, Ga has a concentration of 0.6 mass % or less.
A magnetically anisotropic rare-earth-iron-boron based alloy powder according to the present invention has a mean particle size of 10 μm to 300 μm. The concentration of rare-earth elements in powder particles with sizes of 50 μm or less is not higher than that of rare-earth elements in powder particles of which the sizes exceed 50 μm.
In one preferred embodiment, the powder is decrepitated by a hydrogen process.
A magnetically anisotropic rare-earth-iron-boron based alloy magnet powder according to the present invention includes a rare-earth element at a concentration of 26 mass % to 32 mass %, an α-Fe phase at 5 vol % or less, and Ga at a concentration of 0.6 mass % or less, and includes a fine texture produced by an HDDR process.
A method for producing an anisotropic bonded magnet according to the present invention includes the steps of preparing a magnetically anisotropic magnet powder by any of the methods described above and mixing the magnetically anisotropic magnet powder with a binder and compacting the mixture under an aligning magnetic field.
An anisotropic bonded magnet according to the present invention includes the magnetically anisotropic rare-earth-iron-boron based alloy magnet powder described above.
A motor according to the present invention includes the anisotropic bonded magnet described above.
a) through 1(d) are cross-sectional views schematically illustrating how a master alloy for use to make a magnetically anisotropic magnet powder according to the present invention forms its metal texture.
a) through 2(c) are cross-sectional views schematically illustrating how the metal structure of a master alloy is formed by a strip casting process.
a) through 3(d) are cross-sectional views schematically illustrating how the metal structure of a master alloy is formed by a conventional ingot casting process.
a), 4(b) and 4(c) schematically illustrate the textures of the master alloy of the present invention, a conventional ingot cast alloy, and a strip cast alloy, respectively, at a time T1 before these alloys are subjected to an HDDR process.
a), 5(b) and 5(c) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, at a time T2 after those alloys have started being subjected to the HDDR process (where T1<T2).
a), 6(b) and 6(c) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, at a time T3 after those alloys have started being subjected to the HDDR process (where T2<T3).
a), 7(b) and 7(c) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, at a time T4 after those alloys have started being subjected to the HDDR process (where T3<T4).
a) through 19(e) are schematic representations illustrating how to carry out an HDDR process.
a) is a graph showing relationships between the main phase minor-axis size and the post-HDDR magnetic properties of the master alloy of the present invention, in which the abscissa represents the average main phase minor-axis size.
b) is a graph showing relationships between the rare-earth-rich phase interval and the post-HDDR magnetic properties of the master alloy, in which the abscissa represents the dispersion interval between the rare-earth-rich phases.
a) through 27(e) are cross-sectional views schematically illustrating how a master alloy for use to make a magnetically anisotropic magnet powder according to the present invention forms its metallurgical structure.
The present inventors acquired the basic idea of the present invention by discovering that the metal texture structure of a master alloy to be subjected to an HDDR process had a significant effect on the time it took to carry out a hydrogenation process. The present inventors carried out an HDDR process on master alloys with various texture structures and evaluated the magnetic properties of the resultant HDDR powders. As a result, the present inventors discovered that by using a master alloy having the metal texture shown in
d) schematically illustrates the metal texture of a master alloy for use to make a magnetically anisotropic magnet powder according to the present invention. This master alloy has a structure in which very small rare-earth-rich phases (shown as black dotted regions in
Hereinafter, a preferred method of making the master alloy will be described with reference to
First, as shown in
In an alloy according to the present invention, the first texture layer is not only unnecessary but also harmful as well for the magnetic powder to exhibit huge magnetization after an HDDR process. Still the first texture layer plays important roles because its surface becomes the nucleus of solidification of the second texture layer and controls the cooling rate of the second texture layer. Thus, the first texture layer is an indispensable element to carry out the present invention. The first texture layer preferably accounts for less than 10 vol %, more preferably less than 5 vol %, of the overall alloy. As will be described later, there is a difference in average minor-axis size between the first and second texture layers. Accordingly, if a cross section of the alloy is observed with a microscope, the thickness ratio of these texture layers can be easily calculated, and the volume ratio can be estimated from the thickness ratio.
To form the second texture layer constantly, the solidification rate needs to be controlled strictly. If the solidification rate were too high, the resultant solidified texture would be too fine. However, if the solidification rate was too low, then α-Fe would be produced.
The first texture layer consists essentially of very small R2Fe14B type crystals, of which the crystal grain sizes are less than 20 μm in an average minor-axis size. Also, the crystallographic orientations of the R2Fe14B type crystal grains are not identifiable even when observed optically with a polarizing microscope.
On the other hand, the second texture layer consists essentially of excessively large R2Fe14B type crystals, of which the crystal grain sizes are 20 μm or more in an average minor-axis size and 100 μm or more in an average major-axis size.
With a polarizing microscope, a maze pattern resulting from the crystallographic C planes of the R2Fe14B type compound or a striped pattern parallel to the cooling member is observed in a portion of the second texture layer where the R2Fe14B type compound is present. In the second texture layer, the C-axis of the R2Fe14B type compound is oriented substantially parallel to the cooling member. In other words, this C-axis substantially matches with the minor-axis direction of the crystals.
More specifically, the average minor-axis size of the columnar crystals in the second texture layer is preferably 20 μm to 110 μm, more preferably 60 μm to 110 μm, and most preferably 70 μm to 100 μm. By setting the average minor-axis size of the columnar crystals in the second texture layer within one of these ranges, the coercivity and remanence both increase as will be described later for specific examples of the present invention.
The first and second texture layers, having mutually different average minor-axis sizes, exhibit good magnetic properties when mixed at a predetermined ratio. As described above, if the first texture layer accounts for less than 10 vol % (more preferably less than 5%) of the overall alloy, the HDDR powder exhibits good magnetic properties.
In the R2Fe14B type crystals making up the second texture layer, rare-earth-rich phases are dispersed at an average space of 50 μm or less. The space is preferably 20 μm to 50 μm on average and more preferably 30 μm to 50 μm on average.
The average minor-axis size of the R2Fe14B type crystals in the first and second texture layers is defined herein by the following measuring method. Specifically, a cross section of the alloy as taken in the thickness direction is observed with polarizing micrographs (see
The average minor-axis size is measured herein along the cut lines, which are shifted from the cooling member contact surface in the thickness direction. And the range in which the value is less than 20 μm is defined as the first texture layer, while the range in which the value is 20 μm or more is defined as the second texture layer. Then, the volume percentages described above can be calculated based on the ratio of the thickness of each texture layer to that of the overall alloy.
It should be noted that the average minor-axis size of the R2Fe14B type crystals in the second texture layer is one of the minor-axis sizes to be obtained by the measuring method described above at the center of the thickness of the alloy.
Also, the interval of the rare-earth-rich phases in the second texture layer is obtained by the following measuring method.
When a cross section of the alloy as taken in the thickness direction is observed with backscattering electron images (see
In forming the first texture layer as an aggregation of very small primary crystals, the molten alloy is preferably cooled at a rate of 10° C./s to 1,000° C./s and at a supercooling temperature of 100° C. to 300° C. The supercooling can minimize the nucleation of the Fe primary crystals. On the other hand, in forming the second texture layer, the molten alloy is preferably cooled at a rate of 1° C./s to 500° C./s while being fed continuously.
The cooling rate is adjusted according to the rate of feeding the melt onto the cooling member. Thus, to obtain the alloy texture described above, it is important to adopt a cooling method that allows for adjustment of the melt feeding rate. More specifically, to obtain the alloy texture of the present invention, the melt is preferably constantly fed little by little onto a cooling member (such as a casting mold). For that reason, a cooling process of scattering or atomizing droplets of the melt is preferably carried out. For example, a method of atomizing a melt flow by blowing a gas jet against it or a method of scattering the droplets with centrifugal force may be adopted.
To control the cooling rate of the second texture layer more easily, the following method may be adopted. That method is creating gaps in the first texture layer being formed so as to reduce the substantial heat transfer cross-sectional area of the first texture layer. Then, even if the melt feeding rate is not reduced while the second texture layer is being formed, the cooling rate of the second texture layer decreases as the heat transfer cross-sectional area decreases. Optionally, while the gaps are being created in the first texture layer, the melt feeding rate may be adjusted.
Hereinafter, a preferred method of making the master alloy with those gaps created in the first texture layer will be described with reference to
First, as shown in
By feeding the melt, coarse R2Fe14B type crystals start to grow on the first texture layer as shown in
To create those gaps in the first texture layer, a melt with a relatively high viscosity may be atomized. Specifically, according to a method, the temperature of the melt is decreased from 1,450° C., which is adopted in a normal alloy casting process, to about 1,300° C. or less when the melt reaches the cooling member.
The temperature of the melt may be controlled by a method in which the melt is turned into atomized droplets and then has its heat dissipated while traveling to the cooling member. Specifically, a method in which the atmosphere in a furnace filled with an inert gas is maintained around at the atmospheric pressure or a method of atomizing the melt with an inert gas may be adopted. An Ar gas is normally used as an inert gas. Alternatively, a He gas may be used instead. By using the He gas, the melt droplets can dissipate more heat.
The percentage of the gaps to the overall first texture layer may be represented on the contact surface between the cooling member and the master alloy. If the cross-sectional texture of the master alloy as taken in the thickness direction is observed, then the surface contacting with the cooling member and the gaps can be easily distinguished from each other. Accordingly, the percentage of the gaps may be represented as a ratio of the total length of the gaps to that of the cooling surface. In the alloy of the present invention, the gap percentage falls within the range of 20% to 70%.
Another point in the melt quenching method of the present invention is to collect the produced melt droplets on the cooling member at a high yield (i.e., use the droplets efficiently enough to make a solidified alloy). To increase the yield, a method of blowing the melt droplets onto a flat-plate cooling member with a gas spray or a method of scattering the melt droplets against the inner walls of a rotating drum-like cooling member (i.e., a centrifugal casting process) is preferably adopted.
No solidified alloy with such a texture structure could be obtained by any conventional method such as a strip casting process or an alloy ingot process. Hereinafter, it will be described how crystals grow in a solidified alloy (or master alloy) made by a conventional process.
First, it will be described with reference to
Next, it will be described with reference to
In the master alloy for use in the present invention (see
Using a master alloy with such a texture, even if the concentration of the rare-earth element is defined within the range of 26 mass % to 32 mass %, the α-Fe phase included in the as-cast master alloy yet to be thermally treated can have a small size and its percentage can be reduced to 5 vol % or less. Accordingly, even without the homogenizing heat treatment to be carried out on a master alloy to make a conventional ingot alloy, the magnetic properties of the HDDR powder (e.g., coercivity among other things) would never be affected.
Hereinafter, it will be described what differences will be made when those master alloys having various texture structures are subjected to an HDDR process.
a), 4(b) and 4(a) schematically illustrate the textures of the master alloy of the present invention, a conventional ingot cast alloy, and a strip cast alloy, respectively, at a time T1 before these alloys are subjected to the HDDR process. As shown in
a), 5(b) and 5(a) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, at a time after those alloys have started being subjected to the HDDR process (where T1<T2). In
a), 6(b) and 6(c) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, at a time T3 after those alloys have started being subjected to the HDDR process (where T2<T3 ). As can be seen from
a), 7(b) and 7(c) schematically illustrate the textures of the master alloy of the present invention, the conventional ingot cast alloy, and the strip cast alloy, respectively, during the HDDR process. The textures schematically illustrated in
First, the hydrogenation process was carried out for the amount of time shown in these graphs within a hydrogen atmosphere at a pressure of 0.1 MPa and at a temperature of 850° C. Thereafter, substitution with an argon gas was carried out for 5 minutes. Subsequently, the alloy was subjected to a dehydrogenation process within an argon atmosphere at a temperature of 850° C. and a pressure of 1.0 kPa for 30 minutes. Then, the alloy was cooled to room temperature.
As can be seen from
The remanence Br soon reached its peak after the hydrogenation process had started. Thereafter, the remanence Br decreased as the hydrogenation process lasted longer. This is because the longer the hydrogenation process, the reversible reactions of hydrogenation and dehydrogenation -occur repeatedly, thus gradually erasing the crystallographic orientation that is memorized in the master alloy as described above.
When the alloy of the present invention is used, sufficiently high coercivity HcJ is achieved in a shorter hydrogenation process time than any other alloy. Thus, an HDDR powder with excellent coercivity HcJ and excellent remanence Jr can be obtained.
The coercivity usually decreases as the mean particle size increases. In the alloy of the present invention, however, even if the mean particle size is relatively large, the remanence Jr does not decrease so much. This is believed to be because in the present invention, the master alloy has large crystal grain sizes and a recrystallized texture with aligned crystal orientations should be present in a broader area. In addition, in the present invention, even if the mean particle size increases, the coercivity hardly decreases.
First, master alloys, having the compositions shown in the following Table 1, were made by a centrifugal casting process. More specifically, a melt of a rare-earth-iron-boron based alloy (at a temperature of approximately 1,300° C.) was scattered by centrifugal force toward the inner walls of a rotating cylindrical cooling member. In this manner, master alloys having a texture such as that shown in
Comparing the texture structures of a plurality of alloy samples with different rare-earth contents, the present inventors discovered that the higher the concentration of the rare-earth element included, the smaller the crystal grain size of the alloy.
When a compositional image of coarse crystal grains was observed, it was confirmed that rare-earth-rich phases were dispersed there. The greater the amount of rare-earth elements included in the master alloy, the greater the number of dispersed rare-earth-rich phases identified in the coarse crystal grains. No α-Fe phases were observed.
Next, the master alloys with those various compositions were coarsely pulverized by a hydrogen decrepitation process. More specifically, the alloys were subjected to a hydrogen decrepitation process at 200° C. for 100 minutes within a hydrogen atmosphere, crushed with an agate mortar, and then classified with a sieve, thereby obtaining coarsely pulverized powders with sizes of 425 μm or less.
Thereafter, approximately 10 gram of the coarsely pulverized powder was subjected to an HDDR process. More specifically, a hydrogenation process (in which the temperature rise rate was 15° C./min, the processing temperature was 800° C., the processing time was 1 hour, and a hydrogen atmosphere was used), atmosphere replacement (in which the processing temperature was 800° C., the processing time was 5 minutes, the hydrogen atmosphere was replaced with an argon atmosphere and the flow rate of the argon gas was 5 litters per minute), and a dehydrogenation process (in which the processing temperature was 800° C., the processing time was 1 hour, the argon atmosphere was used continuously and the pressure of the argon gas was 2 kPa) were carried out.
The HDDR processed alloys were classified with a sieve and then the magnetic properties thereof were evaluated with a VSM on a particle size basis. Each sample, along with paraffin, was heated, cooled and fixed under a magnetic field, magnetized with a pulse magnetic field of about 5 MPa, and then the demagnetization curve thereof was plotted.
This particle size dependence of the Nd concentration tends to be opposite to that of the conventional ingot cast alloy or strip cast alloy. That is to say, in the conventional ingot cast alloy or strip cast alloy, the Nd concentration of a fine powder (e.g., with particle sizes of 50 μm or less) is normally higher than that of a coarse powder.
In the conventional ingot cast alloy or strip cast alloy, rare-earth elements such as Nd are present on the grain boundary at a higher concentration than that defined by the stoichiometry of R2Fe14B type crystals and within the crystal grains of the main phase at the value defined by the stoichiometry of R2Fe14B type crystals, respectively. The hydrogen decrepitation process swells the grain boundary portion with the high rare-earth element concentration to make the alloy easy to crack from that portion. Accordingly, the fine powder (with particle sizes of 50 μm or less), included in the coarsely pulverized powder made by the hydrogen decrepitation process, is likely to include the fine powder that has come from the grain boundary. Accordingly, the rare-earth content tends to increase.
In contrast, in the master alloy of the present invention, the rare-earth-rich phases are dispersed within the coarse crystal grains. That's why it is believed that the rare-earth element concentration on the grain boundary is not necessarily higher than that in the main phase in which the rare-earth-rich phases are dispersed. Also, within the main-phase crystal grains of the master alloy, the rare-earth-rich phases are dispersed at an interval of less than about 50 μm (e.g., 10 μm). Accordingly, little rare-earth-rich phases may be included in small powder particles.
Consequently, in the coarsely pulverized powder of the master alloy of the present invention, the concentration of the rare-earth elements included in fine powder particles with mean particle sizes of 50 μm or less is less than that of the rare-earth elements included in powder particles with mean particle sizes exceeding 50 μm. As can be seen from
The alloy that has been subjected to the hydrogen decrepitation and coarse pulverization processes was thermally treated at 800° C. for 1 hour within a vacuum, thereby releasing hydrogen out of the alloy. Thereafter, the size-by-size magnetization of the material powder was measured with a VSM (under an external magnetic field Hex of 1.2 MAm−1).
The magnetization has particle size dependence. That is to say, the greater the particle size, the lower the magnetization tends to be. There is almost no variation in composition among those particle sizes. Accordingly, it is believed that as the particle size increases, the degree of orientation of crystals decreases.
As can be seen from the results shown in
In the prior art, it has been believed preferable to add Co and Ga to the master alloy to obtain an HDDR powder. However, as is clear from these experimental results, even if no Co or Ga is added, a magnet powder with sufficiently good magnetic anisotropy can be obtained according to the present invention. Nevertheless, to reduce the temperature dependence of the magnetic properties, it is effective to add Co. Also, the addition of Co contributes to increasing the anti-corrosiveness. Thus, Co is preferably added according to the intended application. For example, if Co is added to a master alloy including 32 mass % of rare-earth element R, then the content of Co is preferably defined at 1 mass % or more to increase the anti-corrosiveness sufficiently.
In the present invention, if Ga is added, the magnetic properties are improvable to a certain degree as described above. However, Ga is not always indispensable to achieve the objects of the present invention.
By mixing the HDDR powder thus prepared with a known binder and compacting the mixture under a magnetic field, an anisotropic bonded magnet with excellent magnet performance can be obtained. When applied as a permanent magnet to various types of motors and actuators, this anisotropic bonded magnet can exhibit excellent performance.
First, a strip cast alloy and an ingot cast alloy, each having the same composition as Sample No. 10 shown in Table 1, were prepared. Next, each of these alloys was decrepitated by a hydrogenation process so as to be coarsely pulverized to sizes of 425 μm or less. Thereafter, the alloy was subjected to an HDDR process under the following conditions.
After having been once evacuated, the furnace has its internal pressure increased again with an argon gas supplied at the atmospheric pressure (i.e., 0.1 MPa). The samples were heated to 850° C. and then maintained at that temperature with the supply of the argon gas stopped and instead a hydrogen gas started to be supplied. These gases were exhausted (with the pressure maintained and) with an amount of hydrogen gas, corresponding to approximately 20% of the internal volume of the furnace, supplied into the furnace every minute. Such a state was maintained for two hours and then the supply of the hydrogen gas was stopped and instead an argon gas started to be supplied into the furnace with the in-furnace temperature kept substantially constant. In this manner, the argon gas was continuously introduced for five minutes, thereby replacing the atmosphere in the furnace with the argon gas. Thereafter, the in-furnace argon gas pressure was reduced with a rotary pump to 2 kPa, which state was maintained for one hour. Subsequently, an argon gas was supplied into the furnace again, thereby performing a cooling process with the in-furnace argon gas pressure raised to the atmospheric pressure.
This HDDR process is characterized by heating a sample to an elevated temperature (of 550° C. to 900° C.) within a non-hydrogen gas atmosphere and then supplying hydrogen into the furnace to start the hydrogenation process. By delaying the supply of hydrogen into the furnace until the temperature of the alloy has been raised sufficiently, it is possible to prevent the HDDR process from advancing excessively. The master alloy of the present invention is easier to react to hydrogen than the conventional alloy. That is why the HDDR process is preferably somewhat delayed by not allowing the alloy to react to hydrogen until the alloy has been heated to a high temperature.
The powder samples, obtained by the HDDR process described above, were classified with a sieve, and then the remanences Jr and coercivities HcJ of the samples were measured with a VSM on a particle size basis. The results of measurements are shown in
The results shown in
A molten alloy, having a composition consisting essentially of 27.0 mass % of Nd, 1.0 mass % of Dy, 15.0 mass % of Co, 0.6 mass % of Ga, 0.1 mass % of Zr, 1.0 mass % of B and Fe as the balance, was prepared and deposited on a cooling plate by a centrifugal atomization process, thereby making a master alloy. In this case, approximately 50% of gaps were created on its surface contacting with the cooling member. By changing the amount of the melt sprayed, the rate of deposition on the cooling plate was adjusted. In this case, as the amount of the melt sprayed is increased, the deposition rate increases and the cooling rate of the molten alloy decreases. Conversely, as the amount of the melt sprayed is decreased, the deposition rate decreases and the cooling rate of the molten alloy increases. In this manner, master alloys were obtained at various cooling rates.
The cross sections of these master alloys were observed with a microscope, and the particle sizes of the main phases and the dispersion intervals of the rare-earth-rich phases were measured by an image processing technique. More specifically, the dispersion intervals were determined by a cutting process in which cut lines were defined parallel to a cooling substrate.
Then, the master alloy obtained in this manner was not subjected to any particular high-temperature heat treatment and subjected to a hydrogen decrepitation process so as to be coarsely pulverized to sizes of 425 μm or less. Thereafter, the alloy was subjected to an HDDR process, which was carried out in the following manner.
First, the samples were heated up to 800° C. and then maintained at 800° C. for two hours with a hydrogen gas at the atmospheric pressure (i.e., 0.1 MPa) introduced into the furnace. Then the supply of the hydrogen gas was stopped and instead an argon gas started to be supplied into the furnace. In this manner, the argon gas was continuously introduced for five minutes, thereby replacing the atmosphere in the furnace with the argon gas. Thereafter, the in-furnace argon gas pressure was reduced to 1 kPa, which state was maintained for one hour. Subsequently, an argon gas was supplied into the furnace again, thereby performing a cooling process with the in-furnace argon gas pressure raised to the atmospheric pressure. This HDDR process is different from the counterpart of the second example described above in that the samples were heated within the hydrogen gas atmosphere.
a) is a graph showing relationships between the minor-axis size of the main phase and the post-HDDR magnetic properties of the master alloy of the present invention.
Industrial Applicability
According to the present invention, even without adding expensive Ga, an HDDR process can be carried out effectively and a big recrystallized texture with excellent magnetic anisotropy can be produced. As a result, the coercivity HcJ and remanence Jr of the HDDR powder can be both increased. In addition, no homogenizing heat treatment needs to be carried out on the master alloy and the hydrogenation process time of the HDDR process can be shortened. Consequently, the manufacturing cost can be reduced and the manufacturing time can be shortened.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2001-385941 | Dec 2001 | JP | national |
| 2002-030392 | Feb 2002 | JP | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/JP02/13268 | 12/18/2002 | WO | 00 | 5/21/2004 |
| Publishing Document | Publishing Date | Country | Kind |
|---|---|---|---|
| WO03/052779 | 6/26/2003 | WO | A |
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| 20050016632 A1 | Jan 2005 | US |
