R-T-B—Ga-based magnet material alloy and method of producing the same
10497497 · 2019-12-03
Assignee
Inventors
Cpc classification
B22F2301/45
PERFORMING OPERATIONS; TRANSPORTING
B22F2009/048
PERFORMING OPERATIONS; TRANSPORTING
B22F2999/00
PERFORMING OPERATIONS; TRANSPORTING
B22F2009/048
PERFORMING OPERATIONS; TRANSPORTING
C22C30/00
CHEMISTRY; METALLURGY
B22F2998/10
PERFORMING OPERATIONS; TRANSPORTING
B22F2998/10
PERFORMING OPERATIONS; TRANSPORTING
B22F2999/00
PERFORMING OPERATIONS; TRANSPORTING
B22F3/24
PERFORMING OPERATIONS; TRANSPORTING
International classification
B22F9/04
PERFORMING OPERATIONS; TRANSPORTING
Abstract
Disclosed is an R-T-BGa-based magnet material ahoy where R is at least one element selected from rare earth metals including Y and excluding Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu, and T is one or more transition metals with Fe being an essential element. The R-T-BGa-based magnet material alloy includes: an R.sub.2T.sub.14B phase 3 which is a principal phase, and an R-rich phase (1 and 2) which is a phase enriched with the R, wherein a non-crystalline phase 1 in the R-rich phase has a Ga content (mass %) that is higher than a Ga content (mass %) of a crystalline phase 2 in the R-rich phase. With this, it is possible to enhance the magnetic properties of rare earth magnets that are manufactured from the alloy and reduce variations in the magnetic properties thereof.
Claims
1. An R-T-BGa-based magnet material alloy, where R is at least one element selected from rare earth metals including Y and excluding Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu, and T is one or more transition metals with Fe being an essential element, the R-T-BGa-based magnet material alloy, comprising: an R.sub.2T.sub.14B phase which Is a principal phase; and an R-rich phase which is a phase enriched with the R, the R-rich phase Including a non-crystalline phase and a crystalline phase, the non-crystalline phase having a Ga content in mass % that is higher than a Ga content In mass % of the crystalline phase.
2. The R-T-BGa-based magnet material alloy according to claim 1, wherein the R-T-BGa-based magnet material alloy has an average thickness in a range of 0.1 mm to 1.0 mm.
3. A method of producing the R-T-BGa-based magnet material alloy according to claim 1, the method comprising: a first step of casting a ribbon from a molten R-T-BGa-based alloy using a strip casting method and crushing the ribbon to produce alloy flakes; and a second step of thermally maintaining the alloy flakes by holding the alloy flakes at a predetermined temperature for a predetermined time and then cooling the alloy flakes; the first step and the second step being performed under reduced pressure or in an inert gas atmosphere; the second step being performed such that the alloy flakes are thermally maintained at a temperature in a range of 650 C. to the melting temperature of the alloy, and after the thermal maintenance, the alloy flakes are cooled at a cooling rate of 1 C. to 9 C. per second to at least 400 C.
4. A method of producing the R-T-BGa-based magnet material alloy according to claim 2, the method comprising: a first step of casting a ribbon from a molten R-T-BGa-based alloy using a strip casting method and crushing the ribbon to produce alloy flakes; and a second step of thermally maintaining the alloy flakes by holding the alloy flakes at a predetermined temperature for a predetermined time and then cooling the alloy flakes; the first step and the second step being performed under reduced pressure or in an inert gas atmosphere; the second step being performed such that the alloy flakes are thermally maintained at a temperature in a range of 650 C. to the melting temperature of the alloy, and after the thermal maintenance, the alloy flakes are cooled at a cooling rate of 1 C. to 9 C. per second to at least 400 C.
Description
BRIEF DESCRIPTION OF DRAWINGS
(1)
(2)
DESCRIPTION OF EMBODIMENTS
(3) 1. Magnet Material Alloy of the Present Invention
(4) As described above, the magnet material alloy of the present invention is an R-T-BGa-based magnet material alloy where R is at least one element selected from rare earth metals including Y, and T is one or more transition metals with Fe being an essential element, the R-T-BGa-based magnet material alloy comprising: an R.sub.2T.sub.14B phase which is a principal phase and an R-rich phase which is a phase enriched with the R, wherein the R-rich phase includes a non-crystalline phase and a crystalline phase, and the non-crystalline phase has a Ga content that is higher than a Ga content of the crystalline phase. The following provides an explanation of the reasons for specifying the magnet material alloy of the present invention as set forth above and preferred embodiments thereof.
(5) [Alloy Composition]
(6) The magnet material alloy of the present invention is an R-T-BGa-based alloy having a composition that includes, as the R, at least one element selected from rare earth metals including Y; as the T, one or more transition metals with Fe being an essential element; B (boron); and Ga (gallium).
(7) As the R, Nd, Pr, Dy, and Tb are particularly preferred among the rare earth metals including Y, but other rare earth metals such as Sm, La, Ce, Gd, Ho, Er, and Yb may be included. Another embodiment of the invention is one wherein R intentionally does not include or excludes the addition of one or more heavy rare earth elements, that is, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu. Of course, the R-T-BGa based alloy of the invention in any embodiment would include inevitably-contained impurities.
(8) T is one or more transition metals with Fe being an essential element, and Fe alone may constitute the T. The Fe may be partially replaced by Co because Co, among transition metals, produces the effect of enhancing heat resistance. Co acts to reduce the coercive force Hcj in a rare earth-based alloy magnet, but it produces the effect of improving the temperature coefficient of the residual flux density Br. Thus, inclusion of Co increases the squareness of the demagnetization curve, which results in an improved energy product BH (max). Co is preferably included in an amount of 50% or less based on the total T content in order to achieve a balance in magnetic properties for industrial use as a permanent magnet.
(9) The R content preferably ranges from 27.0% by mass to 35.0% by mass. If the R content is less than 27.0% by mass, the amount of rare earth metals necessary for good sintering is not met in the sintering of a green body made from the alloy used as a material for sintered magnets. This results in a decrease in the coercive force. On the other hand, if it exceeds 35.0% by mass, the principal phase becomes relatively small so that the residual flux density Br is decreased. A more preferred R content is from 28.5% by mass to 33.0% by mass although it depends on the desired balance in magnetic properties.
(10) The B content preferably ranges from 0.90% by mass to 1.20% by mass. If it is less than 0.90% by mass, sufficient coercive force Hcj and residual flux density Br may not be achieved in a rare earth magnet made from the alloy. If it exceeds 1.20% by mass, a sufficient residual flux density Br may not be achieved in a rare earth magnet made from the alloy.
(11) [Ga in Non-Crystalline Phase in R-Rich Phase]
(12) The magnet material alloy of the present invention includes an R.sub.2T.sub.14B phase 3 which is a principal phase and an R-rich phase (1 and 2) which is a phase enriched with the R, and the R-rich phase has a non-crystalline phase 1 and a crystalline phase 2 as shown in the later-described Example by referring to
(13) The coercivity mechanism of R-T-B-based sintered magnets is classified as a nucleation type based on the reverse magnetic domain nucleation, and typically the coercive force H.sub.cj can be expressed by the following equation (1).
H.sub.cj=CH.sub.ANI.sub.s(1)
(14) where C is a coefficient representing a decrease in the magnetic anisotropy due to defects near the grain boundaries, surface conditions or the like; H.sub.A is an anisotropy field; N is a demagnetizing factor depending on the size and shape of the grains; and I.sub.s is a saturation magnetization of the principal phase.
(15) As can be seen from the above equation (1), in order to increase the coercive force of an R-T-B-based sintered magnet, it is important to enhance the magnetocrystalline anisotropy H.sub.A of the principal phase as well as to optimize the coefficients C and N, i.e., the balance between the shape, dispersion and the like in the sinter structure.
(16) It is noted that the magnetocrystalline anisotropy H.sub.A of the principal phase is mostly determined by the magnet component system. Thus, optimization of the factors C and N is industrially important. Specifically, increasing the coefficient C and decreasing the coefficient N, i.e., enhancing the interface matching between the principal phase and the R-rich grain boundary phase as well as forming a finer sinter structure leads to an increase in coercive force in an R-T-B-based sintered magnet.
(17) It is known that decreasing the coefficient N, i.e., formation of a finer sinter structure can be achieved to some extent in the sintered magnet manufacturing process. Specifically, the formation of a finer sinter structure can be achieved by making the particle size of the fine powder smaller when the fine powder is prepared by grinding a material alloy, or by lowering the sintering temperature when a green body is sintered.
(18) In the meantime, increasing the coefficient C, i.e., enhancement of the interface matching between the principal phase and the R-rich phase, is greatly affected by the melting point of the R-rich phase. Lowering the melting point of the R-rich phase enhances the interface matching between the principal phase and the R-rich phase. If the melting point of the R-rich phase is lowered, it becomes a melt at an earlier stage of the heating process and the melt viscosity relatively decreases in a conventional temperature range for thermal maintenance (e.g. about 1050 C.) when a green body is sintered in the sintered magnet manufacturing process. Because of this, the wettability between the R-rich phase and the principal phase is improved, which results in enhancing the interface matching.
(19) The magnet material alloy of the present invention has a higher Ga-content non-crystalline phase in its R-rich phase. This is intended to enhance the interface matching between the principal phase and the R-rich phase by lowering the melting point of the R-rich phase. Such non-crystalline phase in the R-rich phase is formed, for example, by thermally maintaining alloy flakes, obtained by crushing a ribbon, under predetermined conditions and then cooling them at a cooling rate of 1 C. to 9 C. per second.
(20) When the slow cooling rate of 1 C. to 9 C. per second is used, a crystalline phase is formed in the R-rich phase while an amorphous phase is formed as nuclei therein because of the presence of Ga in the R-rich phase. The Ga content in the crystalline phase of the R-rich phase is lower than the Ga content in the non-crystalline phase of the R-rich phase. That is, as a greater number of non-crystalline phases, which have a higher Ga content than the crystalline phase, are formed in the R-rich phase, lower melting non-crystalline phases increase. This results in improving the wettability between the R-rich phase and the principal phase to enhance the interface matching when a green body is sintered in the sintered magnet manufacturing process. However, in alloy systems including a heavy rare earth element such as, for example, Dy, Tb, or Ho, if the cooling rate is greater than 9 C. per second, the heavy rare earth element would not sufficiently diffuse to the principal phase in the resulting magnet material alloy. Thus, sintered magnets made from the alloy would tend to exhibit a lower coercive force.
(21) In the meantime, it is known that the residual flux density Br of an R-T-B-based sintered magnet is enhanced as the saturation magnetization I.sub.s of the principal phase increases. Since the saturation magnetization of the principal phase is proportional to the volume of the R.sub.2T.sub.14B phase, which is a ferromagnetic phase, it is necessary to increase the crystallinity, i.e., purity, of the R.sub.2T.sub.14B phase.
(22) The magnet material alloy of the present invention has an R-rich phase that includes a non-crystalline phase having a higher Ga-content. That is, Ga and other impurities are released into the R-rich phase from the principal phase. This is intended to enhance the residual flux density Br of a sintered magnet by increasing the purity of the principal phase in the material alloy. The release of impurities from the principal phase into the R-rich phase is promoted by thermally maintaining the hot alloy flakes, obtained by crushing a ribbon, at a temperature in the range of 650 C. to the melting temperature of the alloy, which activates the elemental diffusion between the R-rich phase and the principal phase. Examples of the impurities that are released from the principal phase into the R-rich phase include Si, Mn and O (oxygen). In particular, it is believed that the melt of Ga, which is non-oxidizing and highly mobile, facilitates the diffusion of the impurities.
(23) As described, the magnet material alloy of the present invention includes a non-crystalline phase and a crystalline phase in its R-rich phase, and the non-crystalline phase in the R-rich phase has a higher Ga content than that of the crystalline phase in the R-rich phase. The higher Ga content non-crystalline phase in the R-rich phase has a lower melting point. This results in improving the wettability between the R-rich phase and the principal phase to enhance the interface matching when a green body is sintered in the process of manufacturing sintered magnets from the alloy of the present invention. Consequently, reverse magnetic domain nucleation is reduced in the resulting sintered magnets and thus the coercive force thereof is enhanced.
(24) Furthermore, when the wettability between the R-rich phase and the principal phase is improved when a green body is sintered in the process of manufacturing sintered magnets from the alloy of the present invention, a homogeneous non-magnetic phase is formed around the principal phase. Because of this, reverse magnetic domain nucleation is reduced and the variations in the coercive force are reduced to provide stabilization in the resulting sintered magnet.
(25) Furthermore, since the magnet material alloy of the present invention is formed to have a higher Ga-content non-crystalline phase in its R-rich phase, the impurity elements within the principal phase are released into the R-rich phase together with Ga, so that the principal phase is cleared and its purity is increased. As such, sintered magnets made from the alloy of the present invention have improved saturation magnetization and an increased residual flux density Br.
(26) The magnet material alloy of the present invention preferably has an average thickness of 0.1 mm to 1.0 mm. It is noted that the average thickness of the magnet material alloy varies in accordance with the thickness of the ribbon that is cast. The average thickness of a magnet material alloy varies from the thickness of a ribbon in accordance with the volume percentage of the R-rich phase, which is a final solidification phase, but the amount of variation is very small. Thus, the average thickness of a magnet material alloy nearly equals the thickness of a ribbon.
(27) If a magnet material alloy having an average thickness smaller than 0.1 mm is to be obtained, the thickness of the ribbon should be smaller than 0.1 mm, too. This would cause excessively rapid cooling of the surface of the ribbon (molten alloy) that contacts with the chill roll and therefore would lead to an increased likelihood of formation of chill grains, which adversely affect the magnetic properties, in the alloy crystal structure. On the other hand, if a magnet material alloy having an average thickness greater than 1.0 mm is to be obtained, the thickness of the ribbon should be greater than 1.0 mm, too. This would lead to a reduced degree of cooling of the ribbon (molten alloy) by the chill roll, which may result in less likelihood of formation of homogeneous columnar dendrites in the alloy crystal structure. In addition, in certain alloy compositions, there may be a problem such as crystallization of an -Fe phase by the peritectic reaction in the alloy crystal structure.
(28) 2. Method of Producing Magnet Material Alloy of the Present Invention
(29) The method of producing a magnet material alloy of the present invention includes producing a magnet material alloy of the present invention as described above, the method comprising: a first step of casting a ribbon from a molten R-T-BGa-based alloy using a strip casting method and crushing the ribbon to produce alloy flakes; and a second step of thermally maintaining the alloy flakes by holding the alloy flakes at a predetermined temperature for a predetermined time and then cooling the alloy flakes; the first step and the second step being performed under reduced pressure or in an inert gas atmosphere, wherein, in the second step, the alloy flakes are thermally maintained at a temperature in a range of 650 C. to the melting temperature of the alloy, and after the thermal maintenance, the alloy flakes are cooled at a cooling rate of 1 C. to 9 C. per second to at least 400 C. The following provides an explanation of the reasons for specifying the method of producing a magnet material alloy of the present invention as set forth above and preferred embodiments thereof.
(30) [First Step]
(31) In the first step, a ribbon is cast from a molten R-T-BGa-based alloy using a strip casting method. Any type of strip casting method may be employed for the ribbon casting as long as it is capable of forming homogeneous columnar dendrites in the crystal structure of the ribbon that is cast by quenching on the contact surface of the chill roll. Thus, either of the following methods may be employed: a single roll method in which a molten alloy is supplied to the outer peripheral surface of a single chill roll; and a twin roll method in which a molten alloy is supplied to a gap formed between two rolls.
(32) When a ribbon is cast using a strip casting method, the ribbon is preferably cast with a thickness of 0.1 to 1.0 mm. When the thickness of the ribbon is smaller than 0.1 mm, the surface of the ribbon (molten alloy) that contacts with the chill roll would be cooled too rapidly. This would lead to an increased likelihood of formation of chill grains, which adversely affect the magnetic properties, in the crystal structure of the cast ribbon. On the other hand, when the thickness of the ribbon is greater than 1.0 mm, the degree of cooling of the ribbon (molten alloy) by the chill roll would be reduced. Thus, problems such as the following occur: a reduced possibility of formation of homogeneous columnar dendrites; or crystallization of an -Fe phase by the peritectic reaction in certain alloy structures.
(33) Alloy flakes are produced by crushing such a ribbon that is cast by a strip casting method.
(34) [Second Step]
(35) In the second step, the alloy flakes produced in the first step as described above are thermally maintained by being held in the hot condition at a predetermined temperature for a predetermined time without being cooled, and thereafter are cooled. During this process, in the method of producing a magnet material alloy of the present invention, the temperature for the thermal maintenance is set to a temperature in the range of 650 C. to the melting temperature of the alloy, and the cooling after the thermal maintenance is performed at a cooling rate of 1 C. to 9 C. per second to at least 400 C.
(36) If the temperature for the thermal maintenance is lower than 650 C., the melting point (eutectic point) of the rare earth-Ga-based intermetallic compound is not reached, and therefore the R-rich phase may not melt to form a liquid phase. On the other hand, if the temperature for the thermal maintenance exceeds the melting temperature of the alloy, part of the alloy melts to adhere to the processing machine. The upper limit of the temperature for the thermal maintenance is preferably 900 C. or less in consideration of, for example, a change in the alloy content due to the formation of the liquid phase.
(37) The hold time for the thermal maintenance is preferably 60 to 1200 seconds although it depends on how much spacing is required between R-rich phases in the magnet material alloy. If the hold time is shorter than 60 seconds, sufficient heating of the liquid phase does not take place, which results in poor elemental diffusion. On the other hand, if the hold time is longer than 1200 seconds, the liquid phase may flow out of the alloy flakes, and this could cause a change in content in the resulting magnet material alloy.
(38) For the cooling after the thermal maintenance, the cooling is performed at a cooling rate of 1 C. to 9 C. per second to at least 400 C. The cooling rate v ( C./second) in the method of producing a magnet material alloy of the present invention may be calculated, for example, by the following equation (2).
v=(T1T2)/t(2)
(39) where t is the time (second) that has elapsed after the start of the cooling; T1 is the temperature ( C.) of the alloy flakes at the start of the cooling; and T2 is the temperature ( C.) of the alloy flakes after the time t has elapsed.
(40) In the method of producing a magnet material alloy of the present invention, cooled to at least 400 C. means: the cooling is performed such that the alloy flakes have a temperature of 400 C. or less at the completion of the cooling, i.e., it is essential that the cooling rate be regulated in the temperature range between the temperature for the thermal maintenance and 400 C. Cooling in a strip casting method often involves slight compositional segregation on the surface of the alloy flakes and inside thereof due to unavoidable inhomogeneity of the surface of the chill roll during the solidification operation or introduction of microoxides or the like during the melting or pouring operation. In the production method of the present invention, the cooling rate is regulated to be in the range of 1 C. to 9 C. per second in the temperature range between the temperature for the thermal maintenance and a sufficiently lower temperature (400 C.) than the liquid phase temperature (about 650 C.). This enables the formation of the higher Ga content non-crystalline phase in the R-rich phase even if there is a slight segregation and presence of a liquid phase having a lower melting point than that of a normal liquid phase.
(41) If the cooling rate for the cooling after the thermal maintenance is slower than 1 C. per second, a sufficient rate of solidification of the R-rich phase melt is not achieved, and therefore it is impossible to form the higher Ga content non-crystalline phase. On the other hand, if the cooling rate is faster than 9 C. per second, Ga, which has an atomic weight smaller, by as much as about 50%, than that of the rare earth metal, a principal component of the R-rich phase, cannot sufficiently diffuse within the R-rich phase, and therefore is impeded from selectively moving to the non-crystalline phase. As a result, the percentage of the lower melting point phase in the non-crystalline phase is decreased. Thus, the wettability between the R-rich phase and the principal phase is not improved when a green body is sintered in the process of manufacturing sintered magnets from the produced alloy, and this results in a decreased coercive force in the manufactured sintered magnets. In addition, in alloy systems including a heavy rare earth element such as, for example, Dy, Tb, or Ho, when the cooling rate is greater than 9 C. per second, sufficient diffusion of the heavy rare earth element to the principal phase does not occur. Because of this, the coercive force is further decreased.
(42) The method of producing a magnet material alloy of the present invention as described includes thermally maintaining alloy flakes, obtained by crushing a ribbon, at 650 C. or higher. This activates the elemental diffusion between the R-rich phase and the principal phase, and therefore promotes the release of impurities including Ga from the principal phase to the R-rich phase. As a result, it is possible to clear the principal phase and increase its purity. In addition, the method of producing a magnet material alloy of the present invention includes cooling the alloy flakes at a cooling rate of 1 C. to 9 C. per second after the thermal maintenance. This allows formation of a crystal structure in which an R-rich phase includes a non-crystalline phase and a crystalline phase and the non-crystalline phase in the R-rich phase has a higher Ga content than that of the crystalline phase in the R-rich phase. As a result, the resulting magnet material alloy has a crystal structure in which its R-rich phase includes a non-crystalline phase and a crystalline phase and the non-crystalline phase in the R-rich phase has a higher Ga content than that of the crystalline phase in the R-rich phase.
EXAMPLES
(43) To verify the advantages of the magnet material alloy of the present invention and the method of producing the same, R-T-BGa-based alloys were prepared and their crystal structures were examined. Also, sintered magnets were made from the prepared R-T-BGa-based alloys, and their magnetic properties were investigated.
(44) 1. Test Method
(45) [Magnet Material Alloy]
(46) In this test, R-T-BGa-based alloy flakes were prepared in accordance with the procedures of Inventive Examples 1 and 2, Conventional Example, and Comparative Examples 1 to 3 as described below. In any of the procedures, the composition of the R-T-BGa-based alloy flakes was as follows: by mass percent, Nd: 24.0%; Pr: 5.0%; Dy: 2.0%; B: 1.0%; and Ga: 0.10% with the balance being Fe and impurities. R-T-BGa-based alloy flakes of this composition have a melting point of about 650 C.
(47) In Inventive Example 1-A, in a chamber under an atmosphere of 300 Torr of argon, 300 kg mass of raw materials were loaded into an alumina crucible and then were melted by radio frequency induction heating to yield a molten alloy. The molten alloy was cast into a ribbon by a strip casting method using a single roll in the chamber. In this process, the molten alloy was supplied to the outer peripheral surface of the chill roll via an alumina tundish. The amount of the molten alloy to be supplied and the number of revolutions of the chill roll were adjusted so that the ribbon was formed to have a thickness of 0.3 mm and thus the resulting alloy flakes had an average thickness of 0.3 mm. The cast ribbon was crushed by a crusher disposed downstream of the chill roll in the chamber to produce alloy flakes.
(48) Subsequently, the produced alloy flakes were loaded into a rotary drum vessel disposed downstream of the crusher in the chamber. In this process, the temperature of the alloy flakes was 762 C. as measured with a two-color pyrometer. The rotary drum vessel used for thermal maintenance and cooling includes a thermal maintenance zone at the upstream side where a heater is disposed and a cooling zone using water cooling at the downstream side. It is thus capable of applying a thermal maintenance treatment and a cooling treatment in sequence to the loaded alloy flakes.
(49) The temperature for the thermal maintenance was adjusted to be 66010 C. and the time for the alloy flakes to pass through the thermal maintenance zone (time for the thermal maintenance) was adjusted to be 613 seconds. This was accomplished by setting the number of revolutions of the rotary drum vessel to 1 rpm and adjusting the output of the heater of the thermal maintenance zone. Thereafter, the alloy flakes were cooled in the cooling zone. In this process, the temperature of the alloy flakes as measured 100 seconds after they entered the cooling zone was 160 C. The cooling rate v for cooling to 160 C. was calculated by the above-described equation (2) using the temperature for the thermal maintenance (660 C.) as the temperature T1 of the alloy flakes at the start of the cooling, and the cooling rate v was found to be 5.0 C. per second.
(50) Subsequently, the alloy flakes discharged from the cooling zone were taken out from the chamber. They were placed in a metallic vessel filled with argon gas, and was allowed to cool in the metallic vessel to room temperature.
(51) In Inventive Example 1, in addition to Inventive Example 1-A as described above in which the alloy flakes had an average thickness of 0.3 mm, Inventive Examples 1-B to 1-D, in each of which the alloy flakes had a different average thickness, were provided. In Inventive Examples 1-B to 1-D, the thickness of the ribbon was varied so that the average thickness of the alloy flakes was varied by adjusting the amount of the molten alloy to be supplied and the number of revolutions of the chill roll. In Inventive Example 1-B, the ribbon was formed to have a thickness of 0.11 mm so that the alloy flakes had an average thickness of 0.11 mm; in Inventive Example 1-C, the ribbon was formed to have a thickness of 0.50 mm so that the alloy flakes had an average thickness of 0.50 mm; and in Inventive Example 1-D, the ribbon was formed to have a thickness of 0.90 mm so that the alloy flakes had an average thickness of 0.90 mm. It is noted that, in Inventive Examples 1-B to 1-D, the cooling rate v was changed with the changes in the thickness of the ribbon and the average thickness of the alloy flakes.
(52) In Inventive Example 2-A, a ribbon was cast by a strip casting method under the same conditions as in Inventive Example 1, and then it was crushed into alloy flakes. In Inventive Example 2, for the thermal maintenance and subsequent cooling performed by loading the crushed alloy flakes into the rotary drum vessel, the temperature for the thermal maintenance was adjusted to be 88010 C. by adjusting the output of the heater of the thermal maintenance zone. The temperature of the alloy flakes when they were to be loaded into the rotary drum vessel was 771 C., and the time for the thermal maintenance was 630 seconds.
(53) Thereafter, when the alloy flakes were cooled in the cooling zone, the temperature of the alloy flakes 100 seconds after they entered the cooling zone was 400 C. The cooling rate v for cooling to 400 C. was calculated by the above-described equation (2) using the temperature for the thermal maintenance (880 C.) as the temperature T1 of the alloy flakes at the start of the cooling, and the cooling rate v was found to be 4.8 C. per second.
(54) In Inventive Example 2, in addition to Inventive Example 2-A as described above in which the alloy flakes had an average thickness of 0.3 mm, Inventive Examples 2-B to 2-D, in each of which the alloy flakes had a different average thickness, were provided. In Inventive Examples 2-B to 2-D, the thickness of the ribbon was varied so that the average thickness of the alloy flakes was varied by adjusting the amount of the molten alloy to be supplied and the number of revolutions of the chill roll. In Inventive Example 2-B, the ribbon was formed to have a thickness of 0.11 mm so that the alloy flakes had an average thickness of 0.11 mm; in Inventive Example 2-C, the ribbon was formed to have a thickness of 0.50 mm so that the alloy flakes had an average thickness of 0.50 mm; and in Inventive Example 2-D, the ribbon was formed to have a thickness of 0.90 mm so that the alloy flakes had an average thickness of 0.90 mm. It is noted that, in Inventive Examples 2-B to 2-D, the cooling rate v was changed with the changes in the thickness of the ribbon and the average thickness of the alloy flakes.
(55) In Conventional Example, an ingot having a thickness of 30 mm and a height of 500 mm was cast from a molten alloy by a mold casting method, and the ingot was crushed to yield alloy flakes.
(56) In Comparative Example 1, a ribbon was cast by a strip casting method under the same conditions as in Inventive Example 1, and then it was crushed into alloy flakes. In Comparative Example 1, for the thermal maintenance and subsequent cooling performed by loading the crushed alloy flakes into the rotary drum vessel, the temperature for the thermal maintenance was adjusted to be 63010 C. by adjusting the output of the heater of the thermal maintenance zone. The temperature of the alloy flakes when they were to be loaded into the rotary drum vessel was 766 C., and the time for the thermal maintenance was 620 seconds.
(57) Thereafter, when the alloy flakes were cooled in the cooling zone, the temperature of the alloy flakes 100 seconds after they entered the cooling zone was 100 C. The cooling rate v for cooling to 100 C. was calculated by the above-described equation (2) using the temperature for the thermal maintenance (630 C.) as the temperature T1 of the alloy flakes at the start of the cooling, and the cooling rate v was found to be 5.3 C. per second.
(58) In Comparative Example 2, a ribbon was cast by a strip casting method under the same conditions as in Inventive Example 1, and then it was crushed into alloy flakes. In Comparative Example 2, for the thermal maintenance and subsequent cooling performed by loading the crushed alloy flakes into the rotary drum vessel, the temperature for the thermal maintenance was adjusted to be 118020 C. by adjusting the output of the heater of the thermal maintenance zone. The temperature of the alloy flakes when they were to be loaded into the rotary drum vessel was 758 C. In Comparative Example 2, a long time, i.e., 920 seconds was required for the alloy flakes to pass through the thermal maintenance zone. Thus, the thermal maintenance zone was inspected, and it was found that a large number of the loaded alloy flakes adhered to the inner surface of the thermal maintenance zone. Because of this, in Comparative Example 2, the test was discontinued with the result that no alloy flakes were produced.
(59) In Comparative Example 3, a ribbon was cast by a strip casting method under the same conditions as in Inventive Example 1, and then it was crushed into alloy flakes. In Comparative Example 3, for the thermal maintenance and subsequent cooling performed by loading the crushed alloy flakes into the rotary drum vessel, the number of revolutions of the rotary drum vessel was varied. Accordingly, the time for the thermal maintenance was 620 seconds. The temperature of the alloy flakes when they were to be loaded into the rotary drum vessel was 766 C.
(60) Thereafter, when the alloy flakes were cooled in the cooling zone, the temperature of the alloy flakes 100 seconds after they entered the cooling zone was 580 C. The cooling rate v for cooling to 580 C. was calculated by the above-described equation (2) using the temperature for the thermal maintenance (660 C.) as the temperature T1 of the alloy flakes at the start of the cooling, and the cooling rate v was found to be 0.8 C. per second.
(61) In Comparative Examples 4 to 7, the average thickness of the alloy flakes was varied in the ribbon casting using a strip casting method. The varying of the average thickness of the alloy flakes was accomplished by changing the thickness of the ribbon by adjusting the amount of the molten alloy to be supplied and the number of revolutions of the chill roll. In Comparative Example 4, the ribbon was formed to have a thickness of 0.08 mm so that the alloy flakes had an average thickness of 0.08 mm; in Comparative Example 5, the ribbon was formed to have a thickness of 0.09 mm so that the alloy flakes had an average thickness of 0.09 mm; in Comparative Example 6, the ribbon was formed to have a thickness of 1.1 mm so that the alloy flakes had an average thickness of 1.1 mm; and in Comparative Example 7, the ribbon was formed to have a thickness of 1.2 mm so that the alloy flakes had an average thickness of 1.2 mm. The other conditions used than these were the same as those of Inventive Example 1, although, in Comparative Examples 4 to 7, the cooling rate v was changed with the changes in the thickness of the ribbon and the average thickness of the alloy flakes.
(62) [Crystal Structure]
(63) Regarding the alloy flakes produced in Inventive Examples 1-A and 2-A, Conventional Example, and Comparative Examples 1 and 3, their crystal structures were examined. In the examination of the crystal structure, for observation of the crystal structure using a transmission electron microscope (TEM), the alloy flakes were treated by ion milling at the surface that contacted with the chill roll and at the surface that was naturally cooled, and specimens were prepared from the thicknesswise central portion thereof. The grain boundaries of the specimens were observed using a transmission electron microscope equipped with a LaB.sub.6 filament at an acceleration voltage of 300 kV.
(64) Also, using energy dispersive X-ray spectroscopy (EDS), an accessory to the transmission electron microscope, the elemental distributions in the principal phase and the R-rich phase of the crystal structure of the alloy flakes were investigated. Furthermore, when it was found that the R-rich phase of the crystal structure includes both a crystalline phase and a non-crystalline phase, elemental distributions in the crystalline phase and the non-crystalline phase of the R-rich phase were also investigated.
(65) From the investigated elemental distributions of the crystalline phase and the non-crystalline phase in the R-rich phase, the Ga contents of the crystalline phase and the non-crystalline phase were each calculated by calculating the mean value of the Ga contents of randomly selected three sites of each phase. Based on the sum of the Ga content of the crystalline phase (mass %) and the Ga content of the non-crystalline phase (mass %) that was calculated, the proportion of the Ga content of the crystalline phase (mass %) and the proportion of the Ga content of the non-crystalline phase (mass %) were each calculated and expressed in percent. These were determined to be the Ga content percentages (%) of the crystalline phase and the non-crystalline phase.
(66) Regarding the alloy flakes produced in Inventive Examples 1 and 2, Conventional Example, and Comparative Examples 1 and 3 to 7, the area percentage of chill grains (%) and the area percentage of -Fe were measured for each of them. In the measurements of the area percentage of chill grains and the area percentage of -Fe, specimens obtained by the following procedure were used.
(67) (1) The produced alloy flakes were taken, and they were mounted in a thermosetting resin and fixed.
(68) (2) For observation of the thicknesswise cross section, the alloy flake fixed in the resin was roughly polished with 120 grit emery paper and then polished by 1200 grit emery paper followed by 3000 grit emery paper, and was mirror finished.
(69) (3) The mirror finished cross section of the alloy flake was etched with nital for five seconds.
(70) Using specimens obtained by the above procedure, the area percentage of chill grains was determined by the following procedure.
(71) (1) Images of the etched cross sections of the alloy flakes were taken at a magnification of 85 using a polarizing microscope.
(72) (2) The taken images were fed into an image processor, and the chill grains were extracted based on fine equiaxed crystal domain.
(73) (3) The area of the chill grains and the cross sectional area of the alloy flake were each calculated, and the area of the chill grains was divided by the cross sectional area of the alloy flake. The result was expressed in percent and defined as the area percentage of chill grains (%).
(74) Also, using specimens obtained by the above procedure, the area percentage of -Fe was determined by the following procedure.
(75) (1) Images of the etched cross sections of the alloy flakes were taken at a magnification of 150 using a scanning electron microscope.
(76) (2) The taken images were fed into the image processor, and the -Fe zone was extracted based on the relative chromaticity (black).
(77) (3) The area of the -Fe zone and the cross sectional area of the alloy flake were each calculated, and the area of the -Fe zone was divided by the cross sectional area of the alloy flake. The result was expressed in percent and defined as the area percentage of -Fe (%).
(78) [Average Thickness of Alloy Flakes]
(79) Regarding the alloy flakes produced in Inventive Examples 1 and 2, Conventional Example, and Comparative Examples 1 and 3 to 7, the average thickness was measured for each of them. In the measurement of the average thickness, ten samples were taken from the produced alloy flakes. Their thicknesses were measured at a central position of the surface of each sample that contacted with the chill roll using a spherical anvil and spindle type micrometer, and the mean value of the thicknesses of the ten samples was calculated.
(80) [Sintered Magnet]
(81) Using the alloy flakes produced in Inventive Examples 1 and 2, Conventional Example, and Comparative Examples 1 and 3 to 7 as a material, sintered magnets were fabricated by the following procedure. Firstly, the alloy flakes were subjected to hydrogen decrepitation (coarse pulverization) by hydrogenation crushing at a hydrogen pressure of 2 kg/cin.sup.2 followed by dehydrogenation in a vacuum at 500 C. for one hour. The coarse powder was milled into a fine powder in a jet mill using a high purity nitrogen at a gas pressure of 6 kg/cm.sup.2. The fine powder had an average particle size of 3.1 m as determined by an air permeability method.
(82) The produced fine powder was pressed at a pressure of 150 MPa in a vertical magnetic field of 2500 kAm.sup.1, thereby being formed into a green body. The green body was sintered at 1050 C. for three hours, and the sintered body was subjected to a heat treatment at 600 C. for one hour to be formed into a permanent magnet.
(83) The heat treated sintered body was cut into pieces sized to 10 mm by 10 mm. Then, they were ground at their end surfaces using a surface grinder to be made into sintered magnets. To determine the residual flux density (Br), energy product ((BH) max) and coercive force (Hcj) of the produced sintered magnets, measurements were made using a BH tracer.
(84) Based on the measurement results, evaluations of the magnetic properties of the sintered magnets were made. In Table 1 below, reference symbols in the evaluation section have the following meanings.
(85) : The symbol indicates good magnetic properties, i.e., a residual flux density Br of 18.0 kG or more, an energy product (BHmax) of 49.0 MGOe or more, and a coercive force (Hcj) of 14.0 kOe or more.
(86) x: The symbol x indicates any of the following is true: a residual flux density Br of less than 18.0 kG, an energy product (BHmax) of less than 49.0 MGOe, and a coercive force (Hcj) of less than 14.0 kOe.
(87) 3. Test Results
(88) Table 1 shows details of each test: the casting method used to produce the R-T-BGa-based alloy; the temperature for thermal maintenance and the cooling rate when alloy flakes are thermally maintained and thereafter cooled; Ga content percentages of the non-crystalline phase and crystalline phase in the R-rich phase of the alloy flakes; and the residual flux density, energy product, coercive force and result of evaluation of magnetic properties of the produced sintered magnet. Table 1 also shows the average thickness of the alloy flakes, the area percentage of chill grains, and the area percentage of -Fe.
(89) [Table 1]
(90) TABLE-US-00001 TABLE 1 Temperature Ga Content Percentage For Thermal In R-Rich Phase Maintenance Cooling Rate Non-Crystalline Crystalline Phase Classification Casting Method ( C.) ( C./sec) phase (%) (%) Inventive Ex. 1-A Strip 660 5.00 88 12 Inventive Ex. 1-B Casting 5.12 87 13 Inventive Ex. 1-C 4.84 90 10 Inventive Ex. 1-D 4.79 89 11 Inventive Ex. 2-A Strip 880 4.80 89 11 Inventive Ex. 2-B Casting 4.92 89 11 Inventive Ex. 2-C 4.49 93 7 Inventive Ex. 2-D 4.35 92 8 Conventional Ex. Mold (Not Detected) 100 Casting Comparative Ex. 1 Strip 630 5.3 41 59 Casting Comparative Ex. 2 Strip 1180 Casting Comparative Ex. 3 Strip 660 0.8 (Not Detected) 100 Casting Comparative Ex. 4 Strip 660 5.5 (Not Examined) (Not examined) Casting Comparative Ex. 5 Strip 660 5.5 (Not Examined) (Not examined) Casting Comparative Ex. 6 Strip 660 3.4 (Not Examined) (Not examined) Casting Comparative Ex. 7 Strip 660 3.2 (Not Examined) (Not examined) Casting Sintered Magnet Area Residual Average Percentage Area Flux Energy Coercive Thickness of Percentage Density Product Force of Alloy Chill Grains of -Fe Classification (kG) (MGOe) (kOe) Evaluation Flakes (mm) (%) (%) Inventive Ex. 1-A 18.1 50.1 14.1 0.30 0 0 Inventive Ex. 1-B 18.1 50.1 14.2 0.11 0 0 Inventive Ex. 1-C 18.5 50.3 14.0 0.50 0 0 Inventive Ex. 1-D 18.4 50.2 14.1 0.90 0 0 Inventive Ex. 2-A 18.0 50.2 14.2 0.30 0 0 Inventive Ex. 2-B 18.0 50.2 14.2 0.11 0 0 Inventive Ex. 2-C 18.4 50.1 14.1 0.50 0 0 Inventive Ex. 2-D 18.1 50.3 14.1 0.90 0 0 Conventional Ex. 14.3 35.1 11.6 x 30 0 0 Comparative Ex. 1 16.5 48.3 13.9 x 0.3 0 0 Comparative Ex. 2 Comparative Ex. 3 17.9 47.5 12.2 x 0.3 0 0 Comparative Ex. 4 16.1 42.3 12.4 x 0.08 5.6 0 Comparative Ex. 5 15.8 41.1 12.1 x 0.09 5.7 0 Comparative Ex. 6 16.2 42.5 12.2 x 1.1 0 2.3 Comparative Ex. 7 16.1 42.1 12.1 x 1.2 0 2.5
(91)
(92)
(93) These results demonstrate that: in the crystal structure of the alloy flakes of Inventive Example 1-A, the non-crystalline phase in the R-rich phase has high contents of O (oxygen), Al, Si, Cu and Ga; the crystalline phase in the R-rich phase has a high content of O (oxygen) while having low contents of Al, Si, Cu and Ga; and the principal phase has low contents of O (oxygen), Al, Si, Cu and Ga.
(94) Regarding the Ga content in the R-rich phase of the alloy flakes of Inventive Example 1-A, it was found that the non-crystalline phase has a higher Ga content than that of the crystalline phase because the Ga content percentage of the non-crystalline phase is higher than that of the crystalline phase. The sintered magnet of Inventive Example 1-A was evaluated as having magnetic properties as represented by the symbol and this confirmed that it has good magnetic properties.
(95) In Inventive Example 2-A, the alloy flakes were thermally maintained at a temperature of 880 C. and thereafter cooled at a cooling rate of 4.8 C. per second. The crystal structure of a specimen prepared from the alloy flakes of Inventive Example 2-A was observed using a transmission electron microscope, and it was found that an R-rich phase was formed at the grain boundaries of a principal phase, and the R-rich phase had a non-crystalline phase and a crystalline phase as is the case with Inventive Example 1-A.
(96) X-ray analyses were performed on the phases of the alloy flakes of Inventive Example 2-A, and it was found that: as is the case with Inventive Example 1-A, the non-crystalline phase in the R-rich phase has high contents of O (oxygen), Al, Si, Cu and Ga; the crystalline phase in the R-rich phase has a high content of O (oxygen) while having low contents of Al, Si, Cu and Ga; and the principal phase has low contents of O (oxygen), Al, Si, Cu and Ga. The sintered magnet of Inventive Example 2-A was evaluated as having magnetic properties as represented by the symbol and this confirmed that it has good magnetic properties.
(97) In Conventional Example, the alloy flakes were produced by crushing an ingot cast by a mold casting method. The crystal structure of a specimen prepared from the alloy flakes of Conventional Example was observed using a transmission electron microscope, and it was found that a principal phase and an R-rich phase were present, but there was no non-crystalline phase in the R-rich phase. X-ray analyses were performed on the phases of the alloy flakes of Conventional Example, and peaks were observed at positions of O (oxygen), Al, Si, Cu, and Ga for both the principal phase and the R-rich phase. The sintered magnet of Conventional Example was evaluated as having magnetic properties as represented by the symbol x, which indicates decreased magnetic properties.
(98) In Comparative Example 1, the alloy flakes were produced by crushing a ribbon cast by a strip casting method, and the alloy flakes were thermally maintained at a temperature of 630 C. and thereafter cooled at a cooling rate of 5.3 C. per second. The crystal structure of a specimen prepared from the alloy flakes of Comparative Example 1 was observed using a transmission electron microscope, and it was found that an R-rich phase was formed at the grain boundaries of a principal phase, and the R-rich phase had a non-crystalline phase and a crystalline phase as is the case with Inventive Example 1.
(99) X-ray analyses were performed on the phases of the alloy flakes of Comparative Example 1, and, unlike the case with Inventive Example 1, peaks were observed at positions of O (oxygen), Al, Si, Cu, and Ga for all the phases, i.e., the non-crystalline phase and the crystalline phase in the R-rich phase, and the principal phase. Furthermore, Table 1 shows that the Ga content percentage of the non-crystalline phase is lower than that of the crystalline phase, and therefore it is seen that, in the R-rich phase, the non-crystalline phase has a lower Ga content than that of the crystalline phase. The sintered magnet of Comparative Example 1 was evaluated as having magnetic properties as represented by the symbol x, which indicates decreased magnetic properties.
(100) In Comparative Example 3, the alloy flakes were thermally maintained at a temperature of 660 C. and thereafter cooled at a cooling rate of 0.8 C. per second. The crystal structure of a specimen prepared from the alloy flakes of Comparative Example 3 was observed using a transmission electron microscope, and it was found that an R-rich phase was present at the grain boundaries of a principal phase, but there was no non-crystalline phase in the R-rich phase.
(101) X-ray analyses were performed on the phases of the alloy flakes of Comparative Example 3, and it was found that the crystalline phase in the R-rich phase has high contents of O (oxygen), Al, Si, Cu and Ga and that the principal phase has low contents of O (oxygen), Al, Si, Cu and Ga. The Ga content of the crystalline phase in the R-rich phase was calculated, and it was found that there was a segregation of Ga in the crystalline phase in the R-rich phase. Also, the sintered magnet of Comparative Example 3 was evaluated as having magnetic properties as represented by the symbol x, which indicates decreased magnetic properties.
(102) These results demonstrate that: by forming an R-T-BGa-based alloy such that it has a crystal structure in which its R-rich phase includes a non-crystalline phase and a crystalline phase and that the non-crystalline phase in the R-rich phase has a higher Ga content than that of the crystalline phase in the R-rich phase, it is possible to provide sintered magnets manufactured from such alloy with enhanced magnetic properties. Also, it has been found that such magnet material alloys can be fabricated by thermally maintaining alloy flakes and thereafter cooling them in such a manner that the temperature for the thermal maintenance is in the range of 650 C. to the melting temperature of the alloy and the cooling rate is 1 C. to 9 C. per second.
(103) In Inventive Examples 1-A to 1-D and Inventive Examples 2-A to 2-D, the alloy flakes were thermally maintained and thereafter cooled in such a manner that the temperature for the thermal maintenance is in the range of 650 C. to the melting temperature of the alloy and the cooling rate is 1 C. to 9 C. per second, and further, the alloy flakes were formed to have an average thickness of 0.1 mm to 1.0 mm. This allowed the R-T-BGa-based alloy to have a crystal structure in which its R-rich phase includes a non-crystalline phase and a crystalline phase, and the non-crystalline phase in the R-rich phase has a higher Ga content than that of the crystalline phase in the R-rich phase. In addition, in the crystal structure of the R-T-BGa-based alloy, the area percentage of chill grains was 0% and the area percentage of -Fe was 0%. That is, there was no formation of chill grains or crystallization of an -Fe phase in the crystal structure of the R-T-BGa-based alloy. Consequently, the sintered magnet was evaluated as having magnetic properties all represented by the symbol , which indicates good magnetic properties.
(104) In contrast, in Comparative Examples 4 and 5, in which the alloy flakes were formed to have an average thickness of less than 0.1 mm, formation of chill grains occurred with their area percentages being 5.6% and 5.7%, respectively, in the crystal structure of the R-T-BGa-based alloy. Consequently, the magnetic properties were evaluated as , which indicates decreased magnetic properties.
(105) In Comparative Examples 6 and 7, in which the alloy flakes were formed to have an average thickness of greater than 1.0 mm, crystallization of an -Fe phase occurred with their area percentages being 2.3% and 2.5%, respectively, in the crystal structure of the R-T-BGa-based alloy. Consequently, the magnetic properties were evaluated as , which indicates decreased magnetic properties.
(106) These results confirmed that, when a ribbon is cast by a strip casting method, the average thickness of the alloy flakes is preferably set to be in the range of 0.1 mm to 1.0 mm.
(107) Although the above embodiment was described with reference to the case, as an example, in which the R-T-BGa-based alloy is used to manufacture a sintered magnet, the present invention is not limited to this. Likewise, when the alloy is used to manufacture bonded magnets, too, it is possible to improve the magnetic properties of resulting bonded magnets.
INDUSTRIAL APPLICABILITY
(108) The magnet material alloy of the present invention has an R-rich phase that includes a non-crystalline phase having a higher Ga-content. Because of this, when it is used as a material for sintered magnets, it is possible to reduce the reverse magnetic domain nucleation in resulting sintered magnets and thus provide improvement and stabilization of the coercive force thereof. Furthermore, it is possible to improve the saturation magnetization and increase the residual flux density of the resulting sintered magnet.
(109) The method of producing a magnet material alloy of the present invention includes thermally maintaining alloy flakes and thereafter cooling the alloy flakes, wherein the temperature for the thermal maintenance is in the range of 650 C. to the melting temperature of the alloy, and the cooling rate is 1 C. to 9 C. per second. This enables the production of a magnet material alloy having a higher Ga-content non-crystalline phase in its R-rich phase.
(110) As described in the foregoing, the magnet material alloy and the method of producing the same of the present invention, when used to manufacture sintered magnets, greatly contribute to improvement in magnetic properties and quality of the resulting sintered magnets. Thus, they can be effectively utilized in the field of rare earth magnets.
REFERENCE SIGNS LIST
(111) 1: non-crystalline phase in R-rich phase,
(112) 2: crystalline phase in R-rich phase,
(113) 3: principal phase