The present invention relates to permanent magnet materials composed mainly of R (where R represents at least one rare earth element), Fe and B, and a process of producing the same, particularly, it relates to R-Fe-B permanent magnet materials and to processes of producing the same. Such materials are typically powdered and then molded into shape and sintered to form magnets.

Nowadays, an R-Fe-B permanent magnet (Japanese Patent Application Laid Open No. Sho 59-46008), is typically used as a high performance permanent magnet. A high magnetic characteristic is obtained by a magnet material structure having a main phase of ternary tetragonal compounds and an R-rich phase, and such magnets are used in a broad field from general domestic electric appliance to peripheral equipment of large-sized computers. Thus R-Fe-B permanent magnets having various structures have been proposed so as to exhibit various magnetic characteristics depending on their proposed uses.

However, in response to recent stringent requirements in the manufacture of small-sized, light and highly functional electric and electronic equipment, inexpensive R-Fe-B permanent magnets with a higher performance are required.

In general, the residual magnetic flux density (Br) of an Fe-B sintered magnet can be expressed as the following Equation (1).$${\text{Br ∞ (Is•β)•f•{ρ/ρ}}_{\text{0}}{\text{•(1-α)}}}^{\text{2/3}}$$ where,

• Is : saturation magnetization
• β : temperature dependence of Is
• f: Degree of orientation
• ρ : density of sintered body
• ρ₀ : theoretical maximum density
• α : volume fraction of grain boundary phase

Thus, in order to raise the residual magnetic flux density (Br) of an R-Fe-B sintered magnet, (1) the volume fraction of the R₂Fe₁₄B matrix phase may be increased, (2) the density of the magnet may be raised to the theoretical maximum density, and further, 3) the degree of orientation of the main phase crystal grains in a easily magnetizing axial direction may be enhanced.

Thus it is important to bring a magnet composition close to the stoichiometrical composition of the above-mentioned R₂Fe₁₄B to achieve the item 1). When the R-Fe-B sintered magnet is produced using as a starting material, an alloy ingot which is prepared by melting the alloy having the aforementioned composition and casting in a mold, since α-Fe crystallizes in the alloy ingot and the R-rich phase becomes segregated locally, it is difficult to pulverize the ingot to fine powders and the composition changes during pulverizing due to oxidation.

In the case of mechanically pulverizing the alloy ingot after hydrogenation and dehydrogenation (as particularly described in Japanese Patent Application Laid Open Nos. Sho 60-63304 and Sho 63-33505) α-Fe crystallized in the alloy ingot remains as it is at the time of pulverization and thus hinders the pulverization because of its ductility, and an R-rich phase which is throughout locally present becomes fine due to hydrogenation producing hydrides. Thus oxidation is accelerated at the time of mechanical pulverization, or pulverization by a jet mill, causing discrepancies in the composition due to dispersion.

When producing the sintered body by using an alloy powder which is brought close to the stoichiometric composition of R₂Fe₁₄B to achieve the item 1), in the sintering process, the presence of a Nd-rich phase for causing liquid phase sintering produces oxides and it is consumed by the inevitable oxidation whereby the sintering is hindered, and since the Nd-rich phase and B-rich phase are inevitably decreased by increase of the R₂Fe₁₄B phase, the production of the sintered body becomes more difficult. Furthermore, the coercive force (iHc) which is one of indexes showing stability of the permanent magnet material and one of its important properties is degraded.

Furthermore, as to the item 3), it is usual in a process of producing an R-Fe-B permanent magnet, to adopt a process of press molding in the magnetic field in order to make the direction of the easy magnetization axes of the main phase crystal grains uniform. In such a case, it is known that the residual magnetic flux density (Br) value and the value of the squareness of the demagnetization curve {(Br2/4(BH)max} change depending on the direction of magnetic field application and the pressing direction, and are influenced by the applied magnetic field intensity.

Recently, a production process has been proposed (Japanese Patent Application Laid Open No. Sho 63-317643) for preventing coarsening of the crystal grains, and residue and segregation of a-Fe which are disadvantages in the production of R-Fe-B alloy powders by an ingot pulverizing process. In that process, a cast piece having a specific thickness is formed from an R-Fe-B molten alloy by the double roll casting method, and according to a common powder metallurgical process, the cast piece is ground coarsely by means of a stamp mill, a jaw crusher or the like, and then comminuted into powders having a mean grain size of 3 to 5 µm by a mechanical pulverizing process in a disk mill a ball mill, grinder, a jet mill or the like, and thereafter pressed in a magnetic field, sintered, and annealed.

However, in this process, as compared with the conventional case of pulverizing an ingot cast in a mold, the pulverizing efficiency at the time of pulverization can not be improved significantly., Also magnetic characteristics can not be greatly improved at the time of pulverization, because not only grain boundary pulverization but also intergranular pulverization occurs, and since the R-rich phase is not in a RH₂ phase stable against oxidation, or since the R-rich phase is fine and has a large surface area, it is poor in oxidation resistance, with the result that oxidation proceeds during the process and high magnetic characteristics can not be obtained.

Recently, demands for cost reduction of the R-Fe-B permanent magnet materials are becoming greater, thus it is very important to manufacture high performance permanent magnet materials efficiently. And hence, manufacturing conditions for drawing out the very best characteristics need to be improved.

We have conducted various studies on processes of producing R-Fe-B permanent magnet materials efficiently and improving their magnetic characteristics.

Enhancement of the residual magnetic flux density (Br) of an R-Fe-B sintered magnet can be achieved by increasing a content of the R₂Fe₁₄B phase of the main phase which is the ferro-magnetic phase. That is, it is important to make the magnet composition close to the stoichiometric composition of R₂Fe₁₄B.

However, when producing the R₂Fe₁₄B sintered magnet from the alloy ingot, prepared by melting the alloy having the aforementioned composition and casting in the mold, as the starting material, particularly as α-Fe crystallized in the alloy ingot and the R-rich phase is locally present throughout, the alloy is difficult to pulverize and has discrepancies in composition.

Also, when producing the alloy powder having the aforementioned composition by a direct reducing and diffusing process, un-reacted Fe grains appear. When raising the reduction temperature to eliminate this, then the grains grow by sintering to one another: if Ca is added as a reducing agent and its oxides are taken in, impurities are thereby increased.

Therefore, as the result of various studies made for the reduction of the various problems related to the production of magnetic alloy materials, we have found out that, by using a strip casting process for rapid cooling and solidifying of the molten alloy, crystallization of an α-Fe phase can be suppressed and a cast alloy piece having a fine grain structure and an homogeneous composition can be produced.

When an R-Fe-B sintered magnet is sintered, a liquid-phase sintering reaction takes place. That is, in the magnet, besides the main R₂Fe₁₄B phase which is the ferromagnetic phase, a B-rich phase and an R-rich phase are present as grain boundary phases, which react with one another during sintering to generate a liquid phase, thereby causing densification of the alloy.

Thus, the B-rich phase and the R-rich phase are indispensable phases for producing a dense R-Fe-B sintered magnet. However, in order to optimise the magnetic characteristics, it is necessary to maximise the R₂Fe₁₄B phase which is the main ferromagnetic phase, and for this purpose, it is necessary to densify an alloy powder which is close to the stoichiometric composition of the R₂Fe₁₄B phase.

It is an object of the present invention to provide high performance R-Fe-B permanent magnet materials having magnetic properties such that when the maximum energy product value (BH)max is expressed in kJ/m³, and the coercive force iHc is expressed in kA/m, the total value$$\frac{\text{(}\mathit{\text{BH}}{\text{)}}_{\text{max}}}{\text{7.96}}\text{+}\frac{\mathit{\text{iHc}}}{\text{79.6}}\text{≥ 59}$$ [this being equivalent to the expression of a total value A + B ≥ 59 wherein (A) is a (BH) max value expressed in megaGaussOersted (1 MGOe = 10⁵ö4π kJ/m³); and (B) is the coercive force iHc value expressed in kilo-Oersted (1 KOe = 10⁶ö4π kA/m)]; and such that when the residual magnetic flux density (Br) is expressed in tenths of one Tesla, (kiloGauss) the squareness of the demagnetization curve {(Br²×1.99)/(BH) max} has a value of between 1.01 and 1.045, [this being equivalent to the expression {(Br²/4)/A} having a value between 1.01 and 1.045 where (A) is again a (BH) max value, but expressed in megaGaussOersted], wherein problems in a process of producing the R-Fe-B materials are solved, efficient pulverization is made possible, oxidation resistance is high, a high iHc is realized by fining crystal grains of a magnet, and orientation of the easy magnetization axis of the crystal grains is improved.

It is another object of the present invention to provide processes of producing such R-Fe-B permanent magnet materials.

According to the present invention, there is provided an R-Fe-B permanent magnet material as defined in Claim 1 hereof.

By the hydrogenation of the strip cast R-Fe-B alloy having a specific composition and thickness, an R-rich phase which is finely dispersed produces hydrides to cause volume expansion and eventual spontaneous decay of the alloy, thereafter the main phase crystal grains constituting the alloy can be comminuted and the powder having a uniform grain distribution can be produced. At this time, the R-rich phase is finely dispersed and the R₂Fe₁₄B phase is also comminuted, thus when the alloy powder which is dehydrogenated and stabilized is comminuted production efficiency is greatly improved, and by orientation using a pulsed magnetic field and pressing, the R-Fe-B permanent magnet of excellent magnetic properties can be obtained.

The present invention also provide various processes for the production of such a magnet material. These processes are defined in Claims 9, 11 and 18 hereof. According to the process of claim 9, an alloy of the desired final composition is directly strip cast and further processed by hydrogenation, dehydrogenation, comminution, magnetic orientation, molding, sintering and annealing. According to the processes of claims 11 and 18, a main phase base alloy and an adjusting alloy are produced separately, and then after hydrogenation, dehydrogenation and comminution the resulting powders are blended together prior to magnetic orientation, molding, sintering and annealing.

For example, by adding and blending an adjusting alloy powder containing a Nd₂Fe₁₇ phase obtained by the strip casting process with the R-Fe-B alloy powder containing the R₂Fe₁₄B phase as the main phase also obtained by the strip casting process, due to the reaction between the Nd₂Fe₁₇ phase in the adjusting alloy powder and the B-rich and Nd-rich phase in the main phase of R-Fe-B alloy powder, the B-rich phase and Nd-rich phase which are deleterious for permanent magnetic characteristics can be adjusted and decreased, the resulting magnet performance can be improved, and further, oxygen content in the alloy powder can be reduced, and an alloy powder having a composition responsive to various magnetic characteristics is provided easily.

Alternatively, by adding and blending an adjusting alloy powder containing an R-Co intermetallic compound phase obtained by the strip casting process with an R-Fe-B alloy powder containing the R₂Fe₁₄B phase as the main phase obtained by the strip casting process, even when the liquid-phase sintering can not be effected only by the main phase of R-Fe-B alloy powder due to the shortage of R-rich and B-rich phases, the R-Co intermetallic compound phase of the adjusting alloy powder is melted to supply a liquid phase for high densification thus the resulting magnet performance can be improved, and further, oxygen content in the alloy powder can be decreased and an alloy powder having a composition responsive to various magnetic characteristics is again provided easily.

Fig. 1 is an explanatory view of a press machine, in which a pulse magnetic field and a static magnetic field can be applied together.

Fig. 2 is a graph showing the relationship between time and magnetic field intensity of a pulse magnetic field.

We have found out that, in the production of an R-Fe-B cast piece having a fine and homogeneous structure by a strip casting process, and comminuting alloy powders which are stabilized by dehydrogenation after hydrogenation, as the result of various studies carried out on a grinding process for the purpose of improving pulverizing efficiency, oxidation resistance and magnetic characteristics of an R-Fe-B sintered magnet, and in particular, the coercive force iHc of an R-Fe-B alloy, pulverizing efficiency is improved about twice as much as conventional pulverizing efficiency, and by molding, sintering and annealing the fine powder which has been oriented by applying a pulse magnetic field, the excellent magnetic properties required by this invention can be achieved.

Thus, when an R-Fe-B alloy which is strip cast and has a specific composition having a structure in which an R-rich phase of specific thickness is finely dispersed, is subjected to hydrogenation of such R-rich phase produces hydrides and causes expansion so that, the alloy can be spontaneously decayed, and as a result, crystal grains constituting the alloy can be comminuted and a powder having a uniform grain distribution can be produced.

It is important that, at this time, the R-rich phase is finely dispersed and the R₂Fe₁₄B phase is also fine. Furthermore, in a process of making the alloy ingot by using a usual mold, when the alloy composition is brought close to the stoichiometric composition of the R₂Fe₁₄B phase, crystallization of α-Fe primary crystals is unavoidable, causing a large deterioration of the pulverizing efficiency in the following process. And hence, though means for providing heat treatment and eliminating α-Fe is taken to homogenize the alloy ingot, since the main phase crystal grains become coarse and segregation of the R-rich phase proceeds, iHc of the sintered magnet is only improved with difficulty.

It is also necessary to make uniform the easily magnetizing axial direction or to improve the degree of orientation of the main phase crystal grains, for achieving high magnetization, and the improvement of the squareness of the demagnetzation curve: and hence, a process of compacting the powder in a magnetic field is adopted.

However, in a coil or a power source disposed on a usual press machine (a hydraulic press or a mechanical press) for generating the magnetic field, a magnetic field of only 10 kOe to 20 kOe (796 to 1592 kA/m) is generated at the most, and the squareness ofthe demagnetization curve {(Br²×1.99)/(BH)ma} also assumes a value of 1.05 or more, thus it is difficult to achieve the theoretical (BH)max value (when the squareness of the demagnetization curve {(Br²×1.99)/(BH)max} is 1.00) expected from a given Br value. Therefore, attempts have been made to mold in a higher magnetic field, but for generating such higher magnetic field, the number of turns of the coil must be increased and also the apparatus requires a high power source and must be made larger.

By analyzing the relationship between the magnetic field intensity at the time of pressing and Br of the sintered body, we have found out that, the more the magnetic field intensity is increased, the higher the magnetization and the more the squareness of the demagnetization curve is improved, thus by using a pulse magnetic field capable of generating a strong magnetic field instantaneously, a higher magnetization and a higher the squareness of the demagnetization curve are made possible.

Meanwhile, we have found out that, in the process of using a pulse magnetic field, it is important instantaneously to orient once by the pulse magnetic field, and it is possible to mold the powder using an isostatic press, and by combining the pulse magnetic field with a static magnetic field generated by an electromagnet, press molding in the magnetic field is also possible.

Thus, according to one aspect of the invention, after strip casting a molten alloy consisting of 12 atomic % to 16 atomic % R (where, R represents at least one rare earth element), 4 atomic % to 8 atomic % B, 5000 ppm or less O₂, Fe (of which a part is optionally replaced by one or both of Co and Ni) and unavoidable impurities, into a cast piece whose main phase is an R₂Fe₁₄B phase, the cast piece is contained in a pressure vessel which can take in and discharge air, the air in the pressure vessel is replaced by hydrogen for hydrogenation of the cast piece. The cast material is then dehydrogenated, and is thereafter comminuted into a fine powder of 1 µm to 10 µm mean particle size under an inert gas, the fine powder is packed into a mold and oriented by applying the pulse magnetic field of 10 kOe (796 kA/m) or more instantaneously, then molded, sintered and annealed, thereby to obtain a permanent magnet material which has magnetic properties such that when the maximum energy product value is expressed in kJ/m³ and the coercive force iHc is measured in kA/m, the total value [(BH)max ö 7.96]+ [iHc ö 79.6] is 59 or more, and such that when the residual magnetic flux density (Br) is measured in tenths of one Tesla (kiloGauss), the squareness of the demagnetization curve {(Br²×1.99)/(BH)max} is 1.01 to 1.045

Whereas an R₂Fe₁₇ phase in an R-Fe alloy such as a Nd-Fe alloy is an intermetallic compound having an easily magnetizing direction in a C phase when its Curie point is in the vicinity of room temperature, and conventionally, in an R-Fe-B sintered permanent magnet, when the amount of B is less than 6 atomic %, for example, an R₂Fe₁₇ phase is produced in the magnet and this weakens its coercivity.

However, as the results of various studies, we have found that, in material powders in which a specific amount of R-Fe alloy powder containing the R₂Fe₁₇ phase such as the Nd₂Fe₁₇ phase is added to and blended with the R-Fe-B alloy powder containing the R₂Fe₁₄B phase as the main phase, near a eutectic temperature of 690°C ofNd in the Nd-rich phase and the Nd₂Fe₁₇ phase in the R-Fe alloy powder in the grain boundary phase, for example, a reaction of Nd + Nd₂Fe₁₇ phase ↔ liquid phase takes place, and this low melting point liquid phase accelerates the sintering of the R-Fe-B alloy powder.

Meanwhile, an adjusting alloy powder containing the Nd₂Fe₁₇ phase and the R-Fe-B alloy powder containing the R₂Fe₁₄B phase as the main phase react as follows during the sintering, and act to increase the R₂Fe₁₄B phase as the main phase. $${\text{13/17 Nd}}_{\text{2}}{\text{Fe}}_{\text{17}}{\text{+ 1/4 Nd}}_{\text{1.1}}{\text{Fe}}_{\text{4}}{\text{B}}_{\text{4}}{\text{+ 133/680 Nd → Nd}}_{\text{2}}{\text{Fe}}_{\text{14}}\text{B}$$

That is, we have found that since the Nd₂Fe₁₄B phase is newly produced by a reaction between the Nd₂Fe₁₇ phase in the adjusting alloy powder and the B-rich phase and Nd-rich phase in the main phase R-Fe-B alloy powder, in a permanent magnet obtained by using only the alloy powder containing the R₂Fe₁₄B phase as the main phase of the conventional process, the amount of the B-rich phase and Nd-rich phase (which is one of the factors deleterious for magnetic characteristics) can be reduced at the time of the sintering reaction.

Furthermore, from the fact that it is a large advantage from a production point of view to obtain material alloy powders which are easily comminuted when producing the R-Fe-B magnet by a powder metallurgical process, as the result of various studies into a process of producing the R-Fe-B permanent magnet material powders, we have found that R-Fe-B permanent magnet material powders may be obtained by mixing a necessary amount of main phase alloy powder containing the R₂Fe₁₄B phase as its main phase and an adjusting alloy powder containing the R₂Fe₁₇ phase also obtained by rapid cooling and solidifying of a molten alloy formed by the strip casting process.

Thus, reasons for producing the main phase alloy powder and adjusting alloy powder from the alloy obtained by the strip casting process in the present invention are that, by strip casting, a main phase alloy powder can be obtained from an alloy cast piece, in which the R₂Fe₁₄B main phase is fine and the B-rich phase and Nd-rich phases are sufficiently dispersed, and in which crystallization of Fe primary crystals is suppressed, and furthermore, an adjusting alloy powder in which the R₂Fe₁₇ phase is dispersed uniformly can be obtained from a strip cast alloy piece.

In particular, when the R₂Fe₁₄B phase is fine and the B-rich phase and R-rich phase are uniformly dispersed in the main phase material powders, pulverizing is improved considerably, and a powder having a uniform particle distribution can be obtained. Furthermore, when producing the magnet, since the crystal structure is fine, a high coercive force is obtained.

Meanwhile, an advantage of producing the adjusting alloy powder containing the R₂Fe₁₇ phase by the strip casting process is that, since the R₂Fe₁₇ phase can be made fine and dispersed sufficiently at the time of mixing with the main phase alloy powder, the reaction takes place uniformly. In the usual alloy melting process using a mold, since α-Fe and the other R-Fe (Co) compound phase are crystallized on the resulting alloy ingot, in order to obtain stable material alloy powders, the alloy ingot must be heated and homogenized, causing the production cost ofthe alloy powder to increase and the R₂Fe₁₇ phase to grow. Furthermore, when producing the adjusting alloy powder by a direct reducing and diffusing process, such problems are encountered that, unreacted Fe grains remain or individual grain compositions differ from each other, and it is very difficult to homogenize the mixture of alloy powders.

As the result of various studies on the above-mentioned findings, we have also found out that, in the material powders prepared by adding and blending a specific amount of R-Co alloy powder containing the R-Co intermetallic compound phase, for example, a Nd₃Co phase or a NdCo₂ phase as the main phase, to an R-Fe-B alloy powder containing the R₂Fe₁₄B phase as the main phase, by the reactions of Nd + Nd₃Co phase ↔ liquid phase in the vicinity of eutectic temperature 625°C ofNd of the Nd-rich phase in the main phase alloy powder and Nd₃Co in the R-Co alloy powder, the low melting point liquid phase accelerates the sintering of the R-Fe-B alloy.

Thus, according to the present invention, it is possible to supply the amount of liquid phase necessary for sintering, as a result, an alloy powder made close to the stoichiometric composition of the R₂Fe₁₄B phase can be liquid-phase sintered. In this way, the magnet composition can be made close to the stoichiometric composition of the R₂Fe₁₄B phase. In other words, when producing the magnet only by the conventional alloy powder containing the R₂Fe₁₄B phase as the main phase, the Nd-rich phase serving as a supply source of the liquid phase produces Nd-oxides during the process by unavoidable material oxidation, so that the amount of liquid phase necessary for sintering can not be secured. As a result, a high sufficiently densification can not be achieved, so that the composition must be set in advance with some tolerance margins. Any such deviations from the optimum composition can be reduced or even eliminated by the present invention.

In particular, when the R₂Fe₁₄B phase in the main phase material powders is fine and the B-rich phase and Nd-rich phase are dispersed uniformly, the comminution is considerably improved at the time of producing the magnet, and a powder having a uniform grain distribution can be produced. Furthermore, since the crystal structure is fine, a high coercivity can be obtained when producing the magnet. Particularly, even when the alloy powder composition is made close to the stoichiometric composition of the R₂Fe₁₄B phase, crystallization of the α-Fe primary crystal is eliminated and a uniform structure is obtained.

Furthermore, advantages of producing the adjusting alloy powder containing the R-Co intermetallic compound phase by the strip casting process are that, such problems as follow can be solved. In the usual alloy melting process using a mold, the Co(Fe) phase and the other R-Co(Fe) compound phase are crystallized in the resulting alloy ingot, and the phases are locally present throughout, therefore, in order to obtain stable material alloy powders, the alloy ingot must be heated and homogenized, causing increase in the production cost of the alloy powder. Also, when producing the adjusting alloy powder by the direct reducing and diffusing process, un-reacted Co and Fe grains remain or individual grain compositions differ from each other, thus it is very difficult to homogenize that whole alloy powder.

It is an object of the present invention to provide high performance R-Fe-B permanent magnet materials having magnetic properties such that

Magnetic characteristics of the R-Fe-B permanent magnet according to the present invention are achieved as follows: when the maximum energy product value (BH)max is expressed in kJ/m³, and the coercive force iHc is expressed in kA/m, the total value$$\frac{\text{(}\mathit{\text{BH}}{\text{)}}_{\text{max}}}{\text{7.96}}\text{+}\frac{\mathit{\text{iHc}}}{\text{79.6}}\text{≥ 59 ;}$$ and also, when the residual magnetic flux density (Br) is expressed in tenths of one Tesla, (kiloGauss) the squareness of the demagnetization curve {(Br²×1.99)/(BH)max} has a value of between 1.01 and 1.045.

This is equivalent to stating that a total value A + B is 59 or more, in which A is a maximum energy product value (BH) max expressed in MGOe and B is a coercive force iHc expressed in kOe, and that the squareness of demagnetization curve {(Br²/4A} value is 1.01 to 1.045 (Br being expressed in tenths of one Tesla [kG]), and A again being a maximum energy product value (BH) max expressed in MGOe.

By selecting the composition and production conditions suitably, the necessary magnetic characteristics are obtained.

In the present invention, a cast piece of the magnet materials having a structure in which the main R₂Fe₁₄B phase and the R-rich phase are finely separated, is produced by strip casting a molten alloy having a specific composition by a single roll process or a double roll process. The resulting cast piece is a sheet whose thickness is 0.03 mm to 10 mm. Though the single roll process and the double roll process may both be used depending on the desired thickness of the cast piece, the double roll process is preferably adopted when the plate thickness is thick, and the single roll process is preferably used when the plate thickness is thin.

Reasons for limiting the thickness of the cast piece within 0.03 mm to 10 mm are that, when the thickness is below 0.03 mm, a rapid cooling effect increases and the crystal grain size becomes smaller than 1 µm, and thus too easily oxidized when comminuted, resulting in deterioration of the magnetic characteristics, and when the thickness exceeds 10 mm, the cooling rate becomes slower, α-Fe is easily crystallized, the crystal grain size becomes larger and also the Nd-rich phase becomes present throughout, and thus the magnetic characteristics deteriorate.

In the present invention, a sectional structure of the R-Fe-B alloy having a given composition obtained by the strip casting process is such that, the main phase R₂Fe₁₄B crystal size is finer than about one tenth or more as compared with that of a conventional ingot obtained by casting in a mold. For example, crystal sizes are 0.1 µm to 50 µm in a short axial direction and 5 µm to 200 µm in a long axial direction, and the R-rich phase is finely dispersed surrounding the main phase crystal grain, and even in local regions, the size is below 20 µm. Crystal grains of the main phase alloy powder and the adjusting alloy powder obtained by the strip casting process have the same properties.

By dispersing the R-rich phase finely below 5 µm, when the R-rich phase produces hydrides at the time of hydrogenation processing, volume expansion occurs uniformly for fracturing, so that the main phase crystal grain is fractured and pulverized and a fine powder having a uniform grain distribution is obtained.

In the following paragraphs, reasons for limiting the compositions of the R-Fe-B permanent magnet and the alloy in the present invention are described.

Rare earth elements which may constitute R contained in the permanent magnet alloy ingot of the present invention may contain yttrium (Y), and are the rare earth elements including light rare earths and heavy rare earths.

As R, the light rare earths are sufficient, and particularly, Nd and Pr are preferable. Though, usually, one kind of R is sufficient, in practice, mixtures (such as mischmetal, didymium, etc.) of two or more rare earth elements will often be used for the reason of availability, and Sm, Y, La, Ce, Gd etc. can be used as a mixture with other R, particularly, Nd, Pr and the like. The R is not necessarily made up of pure rare earth element(s), and those containing unavoidable impurities in production may be used according to what is commercially available.

R is an indispensable element of the alloy for producing the R-Fe-B permanent magnet, in that sufficiently high magnetic characteristics cannot be obtained below 12 atomic %, particularly, a high coercive force can not be obtained, and when exceeding 16 atomic %, the residual magnetic flux density (Br) is lowered and a permanent magnet having the best characteristics can not be obtained. And hence, R is within the range of 12 atomic % to 16 atomic %, the optimum range being 12.5 atomic % to 14 atomic %.

B is an indispensable element of the alloy for producing the R-Fe-B permanent magnet, whereby the high coercive force (iHc) cannot be obtained below 4 atomic %, and when exceeding 8 atomic %, the residual magnetic flux density (Br) is lowered, so that the best permanent magnet cannot be obtained. And hence, the B is 4 atomic % to 8 atomic %, the optimum range being 5.8 atomic % to 7 atomic %.

In the case of Fe, the residual magnetic flux density (Br) is lowered below 76 atomic %, and when exceeding 84 atomic %, the high coercive force cannot be obtained, so that Fe is restricted to 76 to 84 atomic %.

Also, though the reason for substituting a part of Fe with one or both of Co and Ni is to obtain the effect of improved temperature characteristics and corrosion resistance of the permanent magnet, when one or both of Co and Ni exceed 50% of Fe, the highest coercivity cannot be obtained and the best permanent magnet cannot be obtained. And hence, the preferred upper limit of Co and Ni is 50% of Fe.

The reason for restricting O₂ below 5000 ppm is that, when exceeding 5000 ppm, the R-rich phase is oxidized and insufficient liquid phase is produced at sintering, resulting in lowered density, so that a high magnetic flux density cannot be obtained and weathering resistance is also reduced. An optimum range of O₂ is between 200 to 3000 ppm.

When the apparent density of the permanent magnet material is below 7.45 g/cm³, a high magnetic flux density cannot be obtained, and the magnet materials cannot be obtained with a total value [(BH)max ö 7.96]+[iHc ö 79.6] of 59 or more,, which is a feature of the present invention.

Also, as the starting material powders in the present invention, as well as powders of the magnet material composition, it is also possible to blending an R-Fe-B alloy powder containing an R₂Fe₁₄B main phase in which the amount of R, to be described later, is 11 atomic % to 20 atomic %, and an R-Fe-B alloy powder containing the R₂Fe₁₇ phase, in which the amount of R is below 20 atomic % in order to adjust the total amounts of R, B and Fe to the required magnet composition,.

As to the amount of B, the magnet composition can be adjusted by blending the main phase R-Fe-B alloy powder, in which the amount of B is 4 atomic % to 12 atomic % or more, and an adjusting R-Fe-B alloy powder containing the R₂Fe₁₇ phase, in which the amount ofB is below 6 atomic %, or an adjusting R-Fe alloy powder containing the R₂Fe₁₇ phase, in which B is not contained.

Furthermore, the magnet composition can be adjusted by blending an adjusting R-Co (can be substituted by Fe) alloy powder containing an R-Co intermetallic compound (Nd₃-Co, Nd-Co₂ and the like).

Though the presence of unavoidable impurities in industrial production is permissible, besides R, B and Fe in the alloy cast piece of the present invention, by substituting a part of B by a total amount of 4.0 atomic % or less of at least one of 4.0 atomic % or less C, 3.5 atomic % or less P, 2.5 atomic % or less S and 3.5 atomic % or less C, improvements in productivity and reductions in the cost of the magnet alloy are possible.

Meanwhile, by adding at least one of Al―9.5 atomic % or less, Ti―4.5 atomic % or less, V―9.5 atomic % or less, Cr―8.5 atomic % or less, Mn―8.0 atomic % or less, Bi―5 atomic % or less, Nb―12.5 atomic % or less, Ta―10.5 atomic % or less, Mo―9.5 atomic % or less, W―9.5 atomic % or less, Sb―2.5 atomic % or less, Ge―7 atomic % or less, Sn―3.5 atomic % or less, Zr―5.5 atomic % or less and Hf―5.5 atomic % or less, to the alloy powder containing the R, B, Fe alloys or the R-Fe-B alloy containing Co or the blended R₂Fe₁₄B phase as the main phase, or to the adjusting alloy powder containing the R₂Fe₁₇ phase and the adjusting alloy powder containing the R-Co intermetallic compound phase, a high coercivity of the permanent magnet alloy is promoted.

In the R-B-Fe permanent magnet of the present invention, it is necessary that the R₂Fe₁₄B phase of the main phase of a crystal phase presents above 90%, preferably, above 94%. R-Fe-B sintered magnets, which are produced in large numbers at present, has the R₂Fe₁₄B phase of up to 90%. The high magnetic characteristics of the present invention, in which the value [(BH)max ö 7.96]+[iHc ö 79.6] is above 59, can not be obtained below 90%.

A degree of orientation of the magnet of the present invention is calculated from the aforementioned equation 1, it is necessary that the degree of orientation of the magnet is above 85% to hold the value [(BH)max ö 7.96]+[iHc ö 79.6] above 59. When the degree of orientation is below 85%, the squareness of demagnetization curve is poor and the high residual magnetic flux density (Br) is lowered, resulting in a low (BH) max value. The degree of orientation is preferably above 92%.

Though the squareness of the demagnetization curve {(Br²×1.99)/(BH)max} theoretically shows a value of 1.00, since the above-mentioned degree of orientation is inevitably disturbed in the practical permanent magnet material, it has been limited to 1.05 even after many improvement in the past: in the permanent magnet materials of the present invention obtained by the aforementioned specific process, the value of the squareness of demagnetization curve is 1.01 to 1.045.

In the following paragraphs, reasons for the preferred composition ranges of the main phase alloy and the adjusting alloy for the R-Fe-B permanent magnet materials are described.

For obtaining a main phase alloy powder containing the R₂Fe₁₄B phase as the main phase to which an adjusting alloy powder containing the R₂Fe₁₇ phase is added and blended, when R is below 11 atomic %, residual iron where R and B do not diffuse increases, and when exceeding 20 atomic %, the R-rich phase increases and the oxygen content increases at pulverization. Accordingly, R is preferably 11 atomic % to 20 atomic %, more preferably, 13 atomic % to 16 atomic %.

A high coercive force (iHc) can not be obtained when B is below 4 atomic %, and since the residual magnetic flux density (Br) is lowered when exceeding 12 atomic %, the best permanent magnet can not be obtained, so B is preferably 4 atomic % to 12 atomic %, more preferably, 6 atomic % to 10 atomic %.

The rest is composed of Fe and unavoidable impurities, Fe is preferably within the range of 65 atomic % to 82 atomic %. When Fe is below 65 atomic %, the rare earth element(s) and B become relatively abundant, and the R-rich phase and the B-rich phase increase; when exceeding 82 atomic %, the rare earth elements and B decrease relatively, and the residual Fe increases, resulting in a non-uniform alloy powder. Fe is preferably 74 atomic % to 81 atomic %.

When one or both of Co and Ni in the main phase alloy powder are substituted for Fe in the R₂Fe₁₄B main phase this lowers the coercive force, Co is preferably below 10 atomic % and Ni is preferably below 3 atomic %. When replacing a part of Fe with the above-mentioned Co or Ni, Fe is preferably in the range of 55 atomic % to 72 atomic %.

When obtaining an adjusting alloy powder containing the R₂Fe₁₇ phase, the R-rich phase increases in production of the alloy powder and causes oxidation when R exceeds 20 atomic %, thus R is preferably 5 to 15 atomic %. When B is below 6 atomic %, since only the R₂Fe₁₄B phase is present and the amount of B in the main phase alloy powder can be adjusted, B is preferably below 6 atomic %.

Meanwhile, the balance of the main phase powder is composed of Fe and unavoidable impurities, Fe is preferably 85 atomic % to 95 atomic %.

For obtaining an alloy powder containing the R₂Fe₁₄B phase as the main phase, to which the R-Fe adjusting alloy powder containing the R-Co intermetallic compound phase is added and blended, since the residual iron, when R and B do not diffuse, increases when R is below 11 atomic %, and the R-rich phase increases and the oxygen content increases at pulverization when exceeding 15 atomic %, R is preferably 11 atomic % to 15 atomic %, more preferably, 12 atomic % to 14 atomic %.

Since high coercive force (iHc) is not obtained when B is below atomic %, and the residual magnetic flux density (Br) is lowered when exceeding 12 atomic %, B is preferably 4 atomic % to 12 atomic %, more preferably, 6 atomic % to 10 atomic %.

Meanwhile, the balance is composed of Fe and unavoidable impurities, Fe is preferably 73 atomic % to 85 atomic %. When Fe is below 73 atomic %, the rare earth elements and B become abundant relatively and the R-rich phase and the B-rich phase increase, when exceeding 85 atomic %, the rare earth elements and B decrease relatively and the residual Fe increases, results in the non-uniform alloy powder, thus Fe is, more preferably, 76 atomic % to 82 atomic %.

Substituting one or both of Co and Ni in the main phase alloy powder for Fe in the R₂Fe₁₄B main phase tends to reduce the coercive force, so Co is preferably below 10 atomic % and Ni below 3 atomic %. When a part of Fe is replaced with the above-mentioned Co or Ni, Fe is preferably 63 atomic % to 82 atomic %.

For obtaining the adjusting alloy powder containing the R-Co intermetallic compound phase, the R-rich phase increases and tends to cause oxidation in production of the alloy powder when R exceeds 45 atomic %: R is preferably 10 to 20 atomic %.

When the balance is composed of Co and unavoidable impurities, Co is preferably 55 atomic % to 95 atomic %.

One or both of Fe and Ni may be substituted for Co in the adjusting alloy powder. Since the oxidation resistance of the adjusting alloy powder is reduced when the amount of Fe is increased, and the coercive force of the magnet is lowered when the amount ofNi is increased, Fe is preferably below 50 atomic % and Ni below 10 atomic %. When replacing a part of Co with Fe or Ni, Co is preferably 5 atomic % to 45 atomic %.

In the present invention, the magnet composition alloy powder, the main phase alloy powder containing the R₂Fe₁₄B phase as the main phase, and the adjusting alloy powder containing the R₂Fe₁₇ phase or the R-Co intermetallic compound phase, are produced by, for example, a known strip casting process by a single roll process or a double roll process.

Hydrogenation processing is that, for example, a cast piece cut into a predetermined size and having the thickness of 0.03 mm to 10 mm is inserted into a material case, which is covered and charged into a pressure vessel which can be closed tightly, after closing the pressure vessel tightly, the pressure therein is reduced sufficiently, whereafter H₂ gas at 200 Torr (26.6 kPa) to 50 kg/cm² (4.9 MPa) pressure is introduced so that hydrogen is occluded by the cast piece.

Since the hydrogenation reaction is an exothermic reaction, the H₂ gas having a predetermined pressure is supplied for a fixed time, while providing a piping around the pressure vessel for supplying cooling water to suppress the temperature rise in the pressure vessel, so that the H₂ gas is absorbed and the cast piece decays spontaneously and is pulverized. The pulverized alloy is then cooled and dehydrogenated in vacuum.

Since fine cracks are produced in the processed alloy powder grains, it can be comminuted by a ball mill, a jet mill and the like, and the alloy powder having the necessary grain size of 1 µm to 80 µm can be obtained.

In the present invention, air in the processing pressure vessel may be replaced by inert gas beforehand, and with the inert gas being later replaced by the H₂ gas.

The smaller the cut size of the cast piece the lower the H₂ gas pressure required, and though the cut cast piece absorbs H₂ and is comminuted even at low pressures, the higher the hydrogenation pressure, the easier the pulverization. However, the pulverization is reduced when the hydrogen pressure is below 200 Torr (26.6 kPa), and though it may be preferable from a viewpoint of hydrogenation and pulverization to exceed 50 kg/cm² (4.9 MPa), it is not so from the viewpoint of the apparatus and safety, so that the H₂ gas pressure is preferably 200 Torr to 50 kg/cm². For convenience in mass production, it is preferably 2 kg/cm² to 10 kg/cm².

In the present invention, though the pulverization time due to the Hydrogenation varies depending on the closed pressure vessel size, the size of the cut piece and the H₂ gas pressure, it takes more than 5 minutes.

The alloy powder pulverized by hydrogenation is subjected to a primary dehydrogenation in vacuum after cooling. Then, when the pulverized alloy is heated at 100°C to 750°C in vacuum or in argon gas, and subjected to a secondary dehydrogenation for 0.5 hours or longer, the H₂ gas in the pulverized alloy can be completely removed, and oxidation of the powder or a molded body due to prolonged storage is prevented, so that deterioration of the magnetic characteristics of the resulting permanent magnet can be prevented.

Since heating up to 100°C or higher has a good dehydrogenating effect during the dehydrogenation processing, the above-mentioned primary dehydrogenation in vacuum may be omitted, and the decayed powder may be directly dehydrogenated in vacuum or in an argon gas atmosphere at 100°C or higher.

That is, after the hydrogenation and decaying reactions in the aforesaid pressure vessel, the resulting decayed powder may be, subsequently, subjected to dehydrogenation in the pressure vessel atmosphere at 100°C or higher.

Alternatively, after dehydrogenation in vacuum, the decayed powder may be taken out from the pressure vessel for pulverization, whereafter dehydrogenation processing including heating to 100°C or higher in the pressure vessel may be effected again.

When the heating temperature in the above-mentioned dehydrogenation is below 100°C, it takes longer to remove H₂ remaining in the decayed alloy powder thus it is not conducive to mass production. When the temperature exceeds 750°C, a liquid phase is produced and the powder is agglomerated, making pulverization difficult and lessening the moldability at pressing, thus it is not preferable when producing sintered magnets.

When considering the sinterability of the sintered magnet, the preferable dehydrogenation temperature is 200°C to 600°C. Though the processing time varies depending on the processing amount, it usually takes 0.5 hours or longer.

Comminution is suitably effected by a jet mill under an inert gas (e.g. N₂, Ar). It goes without saying that a ball mill or a grinder may be used for comminuting the powder using an organic solvent (e.g. benzene, toluene and the like).

Mean grain sizes ofthe powder at comminution is preferably 1 µm to 10 µm. When below 1 µm, the comminuted powder becomes very active and susceptible to oxidation, with the possibility of spontaneous ignition. When exceeding 10 µm, uncomminuted coarse grains remain to cause deterioration of the coercive force and a slow sintering rate, resulting in a low density. The mean grain size of the fine powder is, more preferably, 2 to 4 µm.

For pressing using the magnetic field, the following process is proposed.

Comminuted powders are packed into a mold in an inert gas atmosphere. The mold may be made of, besides non-magnetic metals and oxides, organic compounds such as plastics, rubber and the like.

The charging density of the powder is from a bulk density (charging density 1.4 g/cm³) in a quiescent state of the powder, up to the solidifying bulk density (charging density 3.0 g/cm³) after tapping. Thus, the charging density is restricted to 1.4 to 3.0 g/cm³.

A pulse magnetic field by an air-core coil and a capacitor power source is applied for orientation of the powder. At the time of orientation, the pulse magnetic field may be applied repeatedly, while compressing the powder by upper and lower punches. The larger the pulse magnetic field intensity, the better, at least 10 kOe (796 kA/m) is necessary, preferably, 30 kOe to 80 kOe (2387 to 6366 kA/m}.

As shown in the graph of Fig. 2 showing the time and the magnetic field intensity, the pulse magnetic field duration is preferably 1 µsec to 10 sec, more preferably 5 µsec to 100 msec, and an applying frequency of the magnetic field is preferably 1 to 10 times, more preferably, 1 to 5 times.

The oriented powder may be solidified by a hydrostatic press. At this time, in the case of using the plastic mold, hydrostatic pressing can be effected as it is. Pressure by the hydrostatic pressing process is preferably 0.5 ton/cm² to 5 ton/cm² (49 to 490 MPa), more preferably, 1 ton/cm² to 3 ton/cm² (98 to 294 MPa).

For continuously performing the orientation by the magnetic field and pressing, it is possible to mold by a magnetic field pressing process, after embedding a coil generating the pulse magnetic field in a die, and using the magnetic field for orientation. Pressure by the magnetic field pressing process is likewise preferably 0.5 ton/cm² to 5 ton/cm² (49 to 490 MPa), more preferably, 1 ton/cm² to 3 ton/cm² (98 to 294 MPa).

A sheet cast piece having a thickness of about 1 mm is prepared from a molten alloy having compositions of Nd 13.0 - B 6.0- Fe 81 obtained by melting in a high frequency melting furnace, by using a double-roll type strip caster including two copper rolls of 200 mm diameter. Crystal grain sizes of the cast piece are 0.5 µm to 15 µm in a short axial direction and 5 µm to 80 µm in a long axial direction. An R-rich phase which is finely separated into about 3 µm is present surrounding a main phase. The oxygen content is 300 ppm.

The cast piece of 1000 g cut into a 50 mm square or smaller is placed in a closed pressure vessel which can take in and discharge air, and N₂ gas is introduced into the pressure vessel for 30 minutes. After flushing out the air, H₂ gas at 3 kg/cm² (about 300 kPa) pressure is fed into the pressure vessel for 2 hours to cause the cast piece to decay spontaneously by hydrogenation. The decayed cast is retained in vacuum at 500°C for 5 hours for dehydrogenation, and thereafter it is cooled to room temperature and ground into 100 mesh.

Next, 800 g of this coarse grain is comminuted in a jet mill to obtain an alloy powder of 3.5 µm mean grain size. The resulting alloy powder is packed into a rubber mold and a pulse magnetic field of 60 kOe (4775 kA/m) is applied instantaneously for orientation. The mold is then subjected to hydrostatic pressing at 2.5 T/cm² (245 MPa) by a hydrostatic press.

The molded body taken out from the mold is sintered at 1090°C for 3 hours to obtain a permanent magnet after one hour annealing at 600°C. Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of demagnetization curve main phase amount and oxygen content are shown in Table 1.

A molten alloy having the same composition as that of Embodiment 1 is strip cast to obtain a sheet cast piece having the sheet thickness of about 0.5 µm

Crystal grain sizes in the cast piece are 0.3 µm to 12 µm in a short axial direction and 5 µm to 70 µm in a long axial direction, and an R-rich phase finely separated into about 3 µm is present surrounding the main phase. The cast piece is comminuted by a jet mill under the same condition as Embodiment 1 to obtain an alloy powder of about 3.4 µm mean grain size. The powder is molded in a magnetic field of about 12 kOe (955 kA/m), after, first, having been oriented in a pulse magnetic field of about 30 kOe (2387 kA/m), by a press machine, in which, as shown in Fig. 1, static magnetic field coils 3,4 are disposed around upper and lower punches 1,2, and a pulse magnetic field coil 6 is provided in a die 5 so as to generate the pulse magnetic field and the usual magnetic field commonly applied to magnet material powders 7. Thereafter, the molded body is sintered and annealed under the same conditions as Embodiment 1.

Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of demagnetization curve, main phase amount and O₂ content of the resulting permanent magnet are shown in Table 1.

An alloy ofNd 13.5 - Dy 0.5 - B 6.5 - Co 1.0 - Fe 78.5 is formed and strip cast as in Embodiment 1 to obtain a sheet cast piece. The cast piece of 100 g cut into a 50 mm square or smaller is decayed spontaneously by hydrogenation as in Embodiment 1, and dehydrogenated in vacuum for 6 hours. Then, after coarse grinding, it is comminuted in a jet mill to obtain powder of 3.5 µm mean grain size.

The resulting powder is oriented in a pulse magnetic field as in Embodiment 1, and a molded body obtained by hydrostatic pressing is sintered similarly. Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of demagnetization curve, main phase amount and O₂ content are shown in Table 1.

The powder obtained at the same condition as the Embodiment 1 is pressed and molded in the magnetic field of about 12 kOe (955 kA/m) by the usual magnetic field press machine in dried state, then sintered and annealed at the same condition as in Embodiment 1. However, oxidation occurred during the pressing, thus densification to a sufficient sinter density was impossible, and the magnetic characteristics could not be measured and only the density and O₂ content are measured (Table 1).

Coarse powder obtained under the same conditions as Embodiment 1 is comminuted in a ball mill, using toluene as a solvent, to obtain the fine powder of 3.5 µm mean grain size, which is pressed and molded in the magnetic field of about 12 kOe (955 kA/m) by the usual magnetic field press machine in a wet state, then sintered and annealed under the same conditions as Embodiment 1.

Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of the demagnetization curve, main phase amount O₂ content of the resulting permanent magnet are shown in Table 1.

A molten alloy having the composition of Nd 14- B 6.0-Fe 80 obtained by melting in a high-frequency melting furnace is cast in an iron mold. When the structure of the resulting alloy ingot was observed, crystallization of a-Fe primary crystals is seen, so it was heated at 1050°C for 10 hours in a homogenizing process.

Crystal grain sizes of the resulting ingot are 30 to 150 µm in a short axial direction and 100 µm to several mm in a long axial direction, and an R-rich phase is segregated with grain sizes of about 150 µm locally.

After coarsely grinding the alloy ingot, the coarse powder is obtained by the hydrogenation and dehydrogenation processes of Embodiment 1. Furthermore, the coarse powder is comminuted by a jet mill under the same conditions as Embodiment 1, and the resulting alloy powder of about 3.7 µm mean grain size is pressed and molded in the magnet field of about 12 kOe (955 kA/m) for sintering and heat treatment at the same conditions as the Embodiment 1. Magnetic characteristics and density, crystal grain size, degree orientation, the squareness of demagnetization curve, main phase amount and O₂ content of the resulting permanent magnet are shown in Table 1.

After coarsely grinding a strip cast piece having the same composition and thickness as the Embodiment 1, 1000 g of the resulting coarse powder is ground, for one hour in a stamp mill, into coarse powders of 100 mesh, without the hydrogenation and dehydrogenation processing, and is then comminuted in the jet mill to obtain the alloy powder of 3.8 µm mean grain size.

The alloy powder is pressed in the magnetic field of about 12 kOe (955 kA/m), sintered and annealed to obtain the permanent magnet. Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of demagnetization curve, main phase amount and O₂ content of the resulting permanent magnet are shown in Table 1.

An alloy having the composition ofNd 13.5- Dy 0.5- B 6.5 - Co 1.0 - Fe 78.5 is cast by the same method as the Comparative Example 3. Since α-Fe primary crystals are present in the resulting alloy ingot, it is subjected to heat treatment at 1050°C for 6 hours. After coarsely grinding the alloy ingot, it is subjected to hydrogenation as in Embodiment 1, and then dehydrogenated in vacuum. The coarse powder is ground coarsely and comminuted in a jet mill to obtain a powder of 3.7 µm mean grain size.

The powder is pressed in the magnetic field of about 12 kOe (955 kA/m), then sintered and heated under the same condition as the Embodiment 1. Magnetic characteristics and density, crystal grain size, degree of orientation, the squareness of demagnetization curve, main phase amount and O₂ content of the resulting permanent magnet are shown in Table 1.

After casting an alloy having the composition ofNd 16.5 - B 7- Fe 76.5 into an ingot as in Comparative Example 3, without liquefaction, the ingot is ground coarsely, and as in Comparative Example 4, coarsely ground in a stamp mill, thereafter comminuted in a jet mill to obtain a fine powder of 3.7 µm mean grain size.

Furthermore, the fine powder is pressed in the magnetic field of about 12 kOe (955 kA/m), then sintered and annealed at the same condition as the Embodiment 1. Magnetic characteristics and density, crystal grain size, degree of orientation the squareness of the demagnetization curve, main phase amount and O₂ content of the resulting permanent magnet are shown in Table 1. Br (kG) Hc (BH)_max iHc 10^-1T kOe kA/m MGOe kJ/m³ kOe kA/m Embodiment 1 14.8 10.50 835.6 53.1 422.6 10.58 841.9 Embodiment 2 14.5 11.0 875.4 50.8 404.3 11.50 915.1 Embodiment 3 13.8 12.9 1026.5 45.9 365.3 15.00 1193.7 Comparative Example 1 _ _ _ _ _ _ _ Comparative Example 2 13.3 9.9 787.8 42.0 334.2 9.98 794.2 Comparative Example 3 13.4 10.3 819.6 42.7 339.8 10.70 851.5 Comparative Example 4 13.1 10.0 795.8 40.5 322.3 10.30 819.6 Comparative Example 5 12.9 11.3 899.2 39.3 312.7 13.50 1074.3 Comparative Example 6 12.2 10.5 835.6 34.4 273.7 11.50 915.1 density ρ (g/cm³) crystal grain size (µm) degree of orientation f(%) angularity (squareness) See Note 1 main phase amount (1-a)(%) oxygen content (ppm) Embodiment 1 7.55 average 6 96 1.031 96.5 1500 Embodiment 2 7.57 average 6 95.5 1.035 94.0 2500 Embodiment 3 7.59 average 6 93.2 1.038 92.7 2000 Comparative Example 1 6.8 _ _ _ _ 6500 Comparative Example 2 7.40 average 11 87.5 1.053 96.5 4200 Comparative Example 3 7.44 average 15 88.4 1.052 95.5 5000 Comparative Example 4 7.43 average 12 86.5 1.060 95.5 5500 Comparative Example 5 7.44 average 14 87.2 1.058 92.7 5000 Comparative Example 6 7.50 average 15 85.8 1.081 86.0 6500 Note 1: The angularity or squareness is calculated as the function {(Br²×1.99)/(BH)max} when Br is measured in tenths of a Tesla and (BH)max is measured in kJ/m³, or as the function {(Br²/4)/(BH)max} when (BH)max is measured in MGOe. These two functions give the same numerical result.

The following are used as materials for a main phase alloy powder by a strip casting process,

• 340 g a Nd metal of 99% purity,
• 8 g of a Dy metal of 99% purity,
• 65.5 g of a Fe-B alloy containing 20% B, and
• 600 g of an electrolytic iron of 99% purity

These materials are melted in an Ar atmosphere so as to obtain an alloy having a predetermined composition, then cast by a strip casting process using copper rolls to obtain a cast piece having a plate thickness of about 2 mm. The cast piece is coarsely ground by a hydrogenation processing, and comminuted by a jaw crusher, a disk mill or the like to obtain 800 g of powder of about 10 µm mean grain size.

The resulting powder consisting of 14.9 atomic % Nd, 0.1 atomic % Pr, 0.3 atomic % Dy, 8.0 atomic % B and Fe, is observed by an x-ray diffraction EPMA, as a result, it is confirmed that O₂ content is about 800 ppm. As the result of EPMA observation on the cast piece structure, the R₂Fe₁₄B main phase is about 5 µm in a short axial direction and 20 to 80 µm in a long axial direction, and the R-rich phase is finely dispersed surrounding the main phase.

As materials of adjusting alloy powders containing an R₂Fe₁₇ phase by the strip casting process,

• 250 g of a Nd metal of 99% purity,
• 11 g of a Dy metal of 99% purity,
• 730 g of an electrolytic iron of 99% purity and
• 20 g of a Fe-B alloy containing 20.0% B

As noted from EPMA observation on the cast piece structure, it consists of the R₂Fe₁₇ phase, partly R₂Fe₁₄B and the Nd-rich phase: α-Fe is not confirmed. The oxygen content is 850 ppm.

Using the above-mentioned two kinds of material powders, 30% adjusting alloy powder is blended with the main phase alloy powder. The material powders are fed into a grinder such as a jet mill or the like to pulverize into about 3 'Lm, the resulting fine powder is filled into a rubber mold, and is subjected to hydrostatic pressing at 2.5 T/cm² (245MPa)by a hydrostatic press machine, after applying a pulse magnetic field of 60 kOe(4775 kA/m) instantaneously for orientation, thereby to obtain a molded body of 8 mm × 15 mm × 10 mm,

The molded body is sintered at 1100°C in an Ar atmosphere for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics ofthe resulting magnet are shown in table 2.

As materials for the main phase alloy powder, as in Embodiment 4,

• 340 g of a Nd metal of 99% purity,
• 8 g of a Dy metal of 99% purity,
• 600 g of an electrolytic iron of 99% purity and
• 65.5 g of a FE-B alloy containing 20% B

As the result of EPMA observation on the alloy ingot structure, the R₂Fe₁₄B main phase is about 50 µm in a short axial direction and about 500 µm in a long axial direction, the R-rich phase having grain sizes of 50 µm is locally present throughout. Also, α-Fe of 5 to 10 µm is seen in the main phase.

As adjusting materials containing the R₂Fe₁₇ phase,

• 200 g Md₂O₃ (98% purity),
• 12 g of Dy₂O₃ (99% purity),
• 65 g of a Fe-B alloy containing 20% B and
• 600 g of iron powders of 99% purity

The molded body is sintered in an Ar atmosphere at 1100°C for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics of the resulting magnet are also shown in Table 2.

The main phase alloy powder of Comparative Example 1 is used, and as materials for the adjusting alloy powder,

• 250 g of a Nd metal of 99% purity,
• 11 g of Dy metal of 99% purity,
• 730 g of an electrolytic iron of 99% purity and 20 g of a Fe-B alloy containing 20.0 g B-are used, melted in an Ar atmosphere and cast in an iron mold. By observation on the structure of the resulting alloy ingot, it is confirmed that a large amount of α-Fe is crystallized, so a homogenizing process is performed at 1000°C for 12 hours.

Component analysis made by the same method as Embodiment 4, shows 10.8 atomic % Nd, 0.1 atomic % Pr, 0.4 atomic % Dy, 2.4 atomic % B and Fe (balance).

Using the above-mentioned two kinds of material powders, 30% adjusting alloy powder is blended with the main phase alloy powder to prepare a magnetic as in Comparative Example 7. Magnetic characteristics of the resulting magnet are shown in Table 2.

As materials,

• 315 g of a Nd metal of 99% purity,
• 8.5 g of a Dy metal of 99% purity,
• 52 g of a Fe-B alloy containing 20 % B and
• 636 g of an electrolytic iron of 99% purity

EPMA observation on the resulting powder, shows 13.8 atomic % Nd, 0.1 atomic % Pr, 0.3 atomic % Dy, 6.3 atomic % B and Fe Balance). The oxygen content is about 800 ppm. As the result EPMA observation also on the cast piece structure, the crystal size ofthe R₂Fe₁₄B main phase is about 6 µm in a short axial direction and 20 to 80 µm in a long axial direction, and the R-rich phase is present as a fine phase surrounding the main phase.

Using the alloy powder by the strip casting process, a magnet is produced as described in Comparative Example 7. Magnetic characteristics of the resulting magnet are also shown in Table 2. composition magnetic characteristics Br (kG) Hc (BH)max iHc 10^-1T kOe kA/m MGOe kJ/m³ kOe kA/m Embodiment 4 13.8Nd - 0.1Pr - 0.3Dy - 6.3B - bal. Fe 14.0 12.5 994.7 47.5 378.0 13.5 1074.3 Comparative Example 7 13.8Nd - 0.1Pr - 0.3Dy - 6.3B - bal. Fe 13.2 12.0 954.9 40.7 323.9 12.5 994.7 Comparative Example 8 13.8Nd - 0.1Pr - 0.3Dy - 6.3B - bal. Fe 13.2 11.9 947.0 40.8 324.7 12.0 954.9 Comparative Example 9 13.8Nd - 0.1Pr - 0.3Dy - 6.3B - bal. Fe 13.3 12.3 978.8 42.3 336.6 12.9 1026.5 density ρ (g/cm³) crystal grain size (µm) degree of orientation f% angularity (squareness) See Note 1 main phase amount (1-α)(%) oxygen content (ppm) Embodiment 4 7.56 average 8 94.5 1.32 92.8 3000 Comparative Example 7 7.53 average 15 89.2 1.07 92.8 5000 Comparative Example 8 7.53 average 16 89.2 1.068 92.8 5500 Comparative Example 9 7.54 average 8 89.9 1.053 92.8 4000 Note 1: The angularity or squareness is calculated as the function {(Br²×1.99)/(BH)max} when Br is measured in tenths of a Tesla and (BH)max is measured in kJ/m³, or as the function {(Br²/4)/(BH)max} when (BH)max is measured in MGOe. These two functions give the same numerical result.

800 g of main phase alloy powder of 10 µm mean grain size having a composition different from the Embodiment 4 is obtained by the process of that Embodiment. The resulting powder consists of 14 atomic % Nd, 0.1 atomic % Pr, 0.5 atomic % Dy, 8 atomic % B and Fe. As the result of observation by the x-ray diffraction EPMA, it is mostly the R₂Fe₁₄B phase. The oxygen content is about 80 ppm. As the result of EPMA observation on the cast piece structure, the R₂Fe₁₄B main phase has a grain size of about 0.5 to 15 µm in a short axial direction and 5 to 90 µm in a long axial direction, and the R-rich phase is dispersed finely to surround the main phase.

As materials ofthe adjusting alloy powder containing an R₂Fe₁₇ phase, 125 g of Nd metal of 99% purity, 5 g of Dy metal of 99% purity and 275 g of an electrolytic iron of 99% purity are used, and a cast piece having a plate thickness of about 2 mm is obtained by the same strip casting process as the main phase alloy. This, powder is prepared by the same process as the main phase alloy. The composition of the resulting powder is 11.0 atomic % Nd, 0.05 atomic % Pr, 0.4 atomic % Dy and Fe.

EPMA observation on the cast piece structure shows it to consist of an R₂Fe₁₇ phase, partly R₂Fe₁₄B, and an R-rich phase: α-Fe is not seen. The oxygen content at 10 µm mean grain size is 700 ppm.

Using the two above-mentioned kinds of material powders, 25 % adjusting alloy powder is blended with the main phase alloy powder. The material powders are charged into a grinder such as a jet mill to comminute it into about 3 µm, then packed into a rubber mold, and the resulting fine powder is subjected to hydrostatic pressing at 2.5 T/cm² (245 MPa) pressure by an isostatic press machine to obtain a molded body of 8 mm x 15 mm x 10 mm, after applying a pulse magnet field of 60 kOe (4775 kA/m) instantaneously for orientation.

The molded body is sintered in an Ar atmosphere at 1100°C for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics of the resulting magnet are shown in Table 3.

As the main phase alloy powder, an alloy having the same composition as the Embodiment 5 is cast in an iron mold to obtain powder of about 10 µm mean grain size by the same method as Embodiment 4. Composition is 14 atomic % Nd, 0.1 atomic % Pr, 0.5 atomic % Dy, 8 atomic % B and Fe (balance), the oxygen content is about 900 ppm. As a result, the crystal grain size is about 50 µm in a short axial direction and about 500 µm in a long axial direction, and an R-rich phase (50 µm) is locally present throughout. Also, 5 to 10 µm crystals of α-Fe are present in the main phase.

The adjusting alloy powder containing the R₂Fe₁₇ phase is produced by the same direct reducing and diffusing process as Comparative Example 7, by using 280 g of Nd₂O₃ (purity 98%), 12 g of Dy₂O₃ (purity 99%) and 750 g of iron powder (purity 99%). Components are 11.0 atomic % Nd, 0.05 atomic % Pr, 0.9 atomic % Dy and Fe. The oxygen content is 1500 ppm.

Using the above-mentioned two kinds of material powder, 25% adjusting alloy powder is blended with the main phase alloy powder, and charged into a jet mill or the like to comminute it to about 3 µm The resulting fine powder is oriented in the magnet field of about 10 kOe (796 kA/m), and molded at about 1.5 T/cm² (about 150 MPa) pressure at right angles to the magnetic field to obtain a molded body of 8 mm × 15 mm × 10 mm.

The molded body is sintered in an Ar atmosphere at 1100°C for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics ofthe resulting magnet are also shown in Table 3.

Using the main phase alloy powder of Comparative Example 10, an adjusting alloy powder is prepared by melting 350 g of a Nd metal, 10 g of a Dy metal and 750 g of an electrolytic iron of 99% purity in the Ar atmosphere, and cast in the iron mold. As the result of observation on the resulting alloy ingot, since a large amount of α-Fe is crystallized, homogenizing is effected at 1000°C for 12 hours. As shown by component analysis, it consists of 11.0 atomic % Nd, 0.05 atomic % Pr, 0.4 atomic % Dy and Fe (balance).

Using the above-mentioned two material powders, 25% adjusting alloy powder is blended with the main phase alloy powder to produce a magnet as in Comparative Example 10. Magnetic characteristics of the resulting magnet are also shown in Table 3.

As starting materials, 300 g of a Nd metal, 13 g of a Dy metal, 50 g of a Fe-B alloy containing 20% B and 645 g of an electrolytic iron of 99% purity are used, and melted in an Ar atmosphere so as to obtain an alloy having a predetermined composition, then, by the strip casting process using copper rolls, a cast piece having a plate thickness of about 2 mm is obtained. Furthermore, the cast piece is pulverized by hydrogenation, and jaw crusher, disk mill and the like to obtain 800 g of powder of about 10 µm mean grain size.

The resulting powder consists of 13.3 atomic % Nd, 0.1 atomic % Pr, 0.5 atomic % Dy, 6 atomic % B and Fe (balance). The oxygen content is about 800 ppm. As shown by EPMA observation on the cast piece structure, the R₂Fe₁₄B main phase crystal size is about 0.3 to 15 µm in a short axial direction and about 5 to 90 µm in a long axial direction, and an R-rich phase is present as a fine phase surrounding the main phase.

Using the alloy powder by the strip casting process, a magnet is produced in the same way as Comparative Example 10. Magnetic characteristics of the resulting magnet are also shown in Table 3. composition magnetic characteristics Br (kG) Hc (BH)max iHc 10^-1T kOe kA/m MGOe kJ/m³ kOe kA/m Embodiment 5 13.3Nd - 0.1Pr - 0.5Dy - 6B - bal. Fe 14.2 12.8 1018.6 48.5 386.0 14.5 1153.9 Comparative Example 10 13.3Nd - 0.1Pr - 0.5Dy - 68 - bal. Fe 13.3 11.5 915.1 41.5 330.2 13.5 1074.3 Comparative Example 11 13.3Nd - 0.1Pr - O.5Dy -6B-bal. Fe 13.3 11.8 939.0 41.7 331.8 13.6 1082.3 Comparative Example 12 13.3Nd -0.1Fr - 0.5Dy - 68 - bal. Fe 13.4 11.6 923.1 42.6 339.0 14.0 1114.1 density ρ (g/cm³) crystal grain size (µm) degree of orientation f% angularity (squareness) See Note 1 main phase amount (1-α)(%) oxygen content (ppm) Embodiment 5 7.57 average 6 95.9 1.039 94.0 2000 Comparative Example 10 7.56 average 14 89.8 1.066 94.0 5000 Comparative Example 11 7.55 average 15 89.8 1.060 94.0 5500 Comparative Example 12 7.56 average 8 90.5 1.054 94.0 3800 Note 1: The angularity or squareness is calculated as the function {(Br²×1.99)/(BH)max} when Br is measured in tenths of a Tesla and (BH)max is measured in kJ/m³, or as the function {(Br²/4)/(BH)max} when (BH)max is measured in MGOe. These two functions give the same numerical result.

As materials of the main phase alloy powder produced by the strip casting process,

• 260 g of a Nd metal of 99% purity,
• 23 g of a Dy metal of 99% purity,
• 68.5 g of a Fe-B alloy containing 20% B and
• 655 g of an electrolytic iron of 99% purity

The resulting powder consists of 11 atomic % Nd, 0.1 atomic % Pr, 1.0 atomic % Dy, 8 atomic % B and Fe (balance), as is observed by x-ray diffraction EPMA, and it is confirmed that it mostly consists of a R₂Fe₁₄B phase. The oxygen content is about 800 ppm. As the result of EPMA observation on the cast piece structure, the R₂Fe₁₄B main phase crystal size is about 0.5 to 1.5 µm in a short axial direction and 5 to 90 µm in a long axial direction, and the R-rich phase is finely dispersed surrounding the main phase.

As materials of an adjusting alloy powder containing an R-Co intermetallic compound phase produced by the strip casting process,

• 490 g of a Nd metal,
• 2∼6 g of a Dy metal and
• 500 g of Co of 99% purity

Following EPMA observation of the cast piece structure, it consists of the R₃Co phase and partly the R₂Co₁₇ phase, and the R₃Co phase is finely dispersed. The oxygen content in the powder, which has a 10 µm mean grain size, is 700 ppm.

Using the above-mentioned two powders, 20% adjusting alloy powder is blended with the main phase alloy powder. The material powders are charged into a grinder such as a jet mill or the like to comminute them into about 3 µm grain sizes, and the resulting powder is filled into a rubber mold and is subjected to hydrostatic pressing at 2.5 T/cm² (245 MPa) by a hydrostatic press machine, after applying a pulse magnetic field of 60 kOe (4775 kA/m) instantaneously for orientation, thereby to obtain a molded body of 8 mm × 15mm × 10 mm.

The molded body is sintered at 1100°C in an Ar atmosphere for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics of the resulting magnet are shown in Table 4.

Magnetic characteristics of the magnet obtained by blending 10 % adjusting alloy powder with the main phase alloy powder prepared in the Embodiment 1, and magnetizing by the same process as the Embodiment 6 are shown in Table 4.

For the main phase alloy powder, as in Embodiment 6,

• 260 g of a Nd metal of 99% purity,
• 26 g of a Dy metal of 99% purity,
• 665 g of an electrolytic iron of 99% purity and
• 68.5 g of a Fe-B alloy containing 20.0% B

EPMA observation on the alloy ingot structure indicates that the R₂Fe₁₄B main phase crystal size is about 50 µm in a short axial direction and about 500 µm in a long axial direction, the R-rich phase (50 µm) is locally present throughout. Some α-Fe of 5 to 10 µm grain size is present in the main phase.

As adjusting materials containing an R-Co intermetallic compound phase, by the direct reducing and diffusing process,

• 550 g of Nd₂O₃ (98% purity),
• 29 g of Dy₂O₃ (99% purity) and
• 500 g of Co powder of 99% purity

Using the above-mentioned two kinds of material powders, 20 % adjusting alloy powder is blended with the main phase alloy powder, and charged into a grinder such as a jet mill or the like to pulverize into about 3 µm. The resulting fine powder is oriented in the magnetic field of about 10 kOe (796 kA/m), and molded at about 1.5 T/cm² (about 150 MPa) pressure to obtain a molded body of 8 mm × 15 mm × 10 mm.

The molded body is sintered in an Ar atmosphere at 1100°C for 3 hours, and annealed at 550°C for one hour. Magnetic characteristics of the resulting magnet are also shown in Table 4.

Using the main phase alloy of Embodiment 13, the adjusting alloy powder is prepared by melting.

• 490 g of a Nd metal,
• 26 g of Dy metal and
• 500 g of Co of 99% purity

Using the above-mentioned two material powders, 20% adjusting alloy powder is blended with the main phase alloy powder to produce a magnet as in Comparative Example 13. Magnetic characteristics of the resulting magnet are also shown in Table 4.

• 305 g of a Nd metal,
• 26 g of a Dy metal,
• 55 g of a Fe-B alloy containing 20% B,
• 100 g of Co of 99% purity, and
• 525 g of an electrolytic iron of 99% purity are melted in an Ar atmosphere so as to obtain an alloy having a predetermined composition, and by the strip casting process using copper rolls, a cast piece having a plate thickness of about 2 mm is obtained. The cast piece is coarsely ground by hydrogenation processing and comminuted by a jaw crusher, disk mill or the like to obtain 800 g of powder of about 10 µm grain size.

The resulting powder consists of 13.5 atomic % Nd, 0.1 atomic % Pr, 1.0 atomic % Dy, 6.7 atomic % B, 11.3 atomic % Co and Fe. The oxygen content is about 800 ppm. As the result of EPMA observation on the cast piece structure, the crystal size of the R₂(Fe, Co₁₄)B phase is about 0.3 to 1.5 µm in a short axial direction and about 5 to 90 µm in a long axial direction, the R-rich phase and the R-Co phase being present finely surrounding the main phase.

Using the alloy powder by the strip casting process, a magnet is produced as in Comparative Example 3.

Magnetic characteristics of the resulting magnet are also shown in Table 4. composition magnetic characteristics Density Br Hc (BH)max iHc 10^-1T kOe kA/m MGOe kJ/m³ kOe kA/m g/cm³ Embodiment 6 13.5Nd - 0.1Pr - 1.ODy 6.7B - 6.5Co - bal. Fe 13.3 12.4 986.8 42.5 338.2 17.0 1352.8 7.62 Embodiment 7 12.3Nd-0.1Pr - 1ODy 7.3B - 11.3Co - bal. Fe 13.5 12.5 994.7 44.0 350.1 16.8 1336.9 7.61 Comparative Example 13 13.SNd - 0.1Pr - 1ODy 6.7B - 11.3Co - bal. Fe 12.0 11.0 875.4 34.0 270.6 15.8 1257.3 7.56 Comparative Example 14 13.5Nd - 0.1Fr - 1ODy 6.7B - 11.3Co - bal. Fe 12.2 11.1 883.3 35.0 278.5 15.5 1233.5 7.55 Comparative Example 15 13.5Nd - 0.1Pr - 1ODy 6.7B - 11.3Co - bal. Fe 12.2 11.2 891.3 35.2 280.1 16.5 1313.0 7.58 density ρ (g/cm³) crystal grain size (µm) degree of orientation f% angularity (squareness) See Note 1 main phase amount (1-α)(%) oxygen content (ppm) Embodiment 6 7.62 average 5 94 1.04 91 2800 Embodiment 7 7.61 average 6 95.5 1.036 94 2200 Comparative Example 13 7.56 average 14 85.7 1.056 91 4800 Comparative Example 14 7.55 average 15 87.1 1.063 91 5000 Comparative Example 15 7.58 average 6 87.1 1.057 91 3500 Note 1: The angularity or squareness is calculated as the function {(Br²×1.99)/(BH)max} when Br is measured in tenths of a Tesla and (BH)max is measured in kJ/m³, or as the function {(Br²/4)/(BH)max} when (BH)max is measured in MGOe. These two functions give the same numerical result.