EP0599815A1

EP0599815A1 - Magnetic alloy and method of making the same

Info

Publication number: EP0599815A1
Application number: EP94101456A
Authority: EP
Inventors: Osamu C/O Seiko Epson Corporation Kobayashi; Koji Osamu C/O Seiko Epson Corporation Akioka; Tatsuya Osamu Seiko Epson Corporation Shimoda
Original assignee: Seiko Epson Corp
Current assignee: Seiko Epson Corp
Priority date: 1987-04-30
Filing date: 1987-09-28
Publication date: 1994-06-01
Anticipated expiration: 2007-09-28
Also published as: EP0599815B1

Abstract

A rare earth-iron magnetic alloy comprising an alloy of between about 8 to 30 atomic percent of at least one rare earth element, between about 2 and 8 atomic percent boron and the balance iron, wherein said magnetic alloy has a columnar macrostructure and an easy axis of magnetisation aligned in a plane perpendicular to the direction of crystal growth. A method of manufacturing a rare earth-iron magnetic alloy is also disclosed; comprising casting a molten alloy composition including at least one rare earth element, iron and boron to form an anisotropic cast ingot having a columnar macrostructure.

Description

This invention relates to magnetic alloys and methods of making the same.
Magnetic alloys, and especially permanent magnets made therefrom, are used in a wide field from household electric appliances to peripheral console units of large sized computers.
With a recent growing demand for smaller sized and higher efficiency electrical appliances, permanent magnets have increasingly been required to have a higher and higher performance. Typical permanent magnets now in use are alnico, hard ferrite magnets and rare earth element - transition metal magnets. In particular, a R-Co (R represents a rare earth element) permanent magnet and a R-Fe-B permanent magnet, which are rare earth element-transition metal magnets, can produce a high magnetic performance, so that much research has hitherto been conducted in relation to them.
For example, the following methods have been developed for manufacturing an R-Fe-B permanent magnet:

(1) Sintering an R-Fe-B material using a powder metallurgy technique (References 1 and 2).
(2) A rapidly quenched thin fragment of about 30 µm (microns) thickness is produced by a melt spinning apparatus which is used for manufacturing an amorphous alloy. A magnet is produced from the thin fragment by resin bonding (References 3 and 4).
(3) Mechanically aligning a rapidly quenched thin fragment obtained by method 2 by a two-stage hot pressing method (References 4 and 5).

Reference 1: Japanese Patent Laid-Open No. 46008/1984.
Reference 2: M. Sagawa, S. Fujimura, N. Togawa, H. Yamamoto and Y. Matushita: J. Apl. Phys. Vol, 55 (6), 15 March, p. 2083 (1984).
Reference 3: Japanese Patent Laid-Open No. 211549/1983.
Reference 4: R.W. Lee: Appl. Phys. Lett. Vol 46 (8), 15 April, p. 790 (1985).
Reference 5: Japanese Patent Laid-Open No. 100402/1985.
In method (1), an alloy ingot is first made by melting and casting, and pulverised to obtain a metal powder having an appropriate particle size (several microns). Powder is needed with a binder of a moulding additive, and pressed in a magnetic field to obtain a moulded body. The moulded body is sintered at approximately 1100^oC in an argon atmosphere for 1 hour and thereafter rapidly quenched to room temperature. After sintering, the sintered body is heat treated at approximately 600^oC to improve the coercive force.
In method (2) a rapidly quenched thin fragment of an R-Fe-B alloy is first made using a melt spinning apparatus at an optimum substrate velocity. The thus obtained ribbon-like thin fragment having a thickness of 30 µm microns) is an aggregate of crystals having a diameter of not more than 100 µm (1,000 Angstroms). It is brittle and easily broken. Since the crystal grains are distributed isotropically, the thin fragment is magnetically isotropic. The thin fragment is pulverised to an appropriate particle size and kneaded with a resin. The mixture is then pressed.
Method (3) produces a dense R-Fe-B magnet having anisotropy by subjecting the rapidly quenched thin fragment obtained by method (2) to a so-called two-stage hot pressing process in vacuum or in an inert gas atmosphere.
During the hot pressing process, uniaxial pressure is applied so as to align the easy magnetisation axis in parallel to the pressing direction and to make the magnet anisotropic.
It is so designed that the crystal grains of the thin fragment produced by the melt spinning apparatus have a smaller grain diameter than the grain diameter of the crystal grains which exhibit the largest coercive force so that the optimum grain diameter is obtained when the crystal grains are made coarser during the subsequent hot pressing process.
Although the above described prior art methods can be used to manufacture permanent magnets containing a rare earth element, iron and boron as the main ingredients they have the following defects.
Method (1) necessitates the step of powdering an alloy. Since an R-Fe-B alloy is very active in oxygen, the step of powdering further accelerates oxidation, whereby the oxygen concentration of the sintered body is inconveniently raised. Furthermore, when the powder is moulded, a moulding additive such as zinc stearate must be used. The moulding additive is removed from the moulded body prior to the sintering step, but several percent thereof remains in the magnet in the form of carbon. This carbon unfavourably significantly lowers the magnetic performance of the R-Fe-B alloy.
The moulded body obtained by press moulding the alloy with the moulding additive added thereto is called a green body, which is very brittle and difficult to handle. It therefore disadvantageously requires much labour to insert the green bodies in a sintering oven in a neatly arranged form. Thus, manufacture of an R-Fe-B permanent magnet generally not only requires expensive equipment, but also has low productivity, resulting in high manufacturing costs. It cannot, therefore, be said that method (1) is capable of making the best use of the merit of an R-Fe-B magnet which is comparatively inexpensive in material costs.
Both methods (2) and (3) use a vacuum melt spinning apparatus. This apparatus has very low productivity and is very expensive at present.
Method (2) adopting a resin bonding process produces a magnet having theoretical isotropy and, hence, a low energy product. Since the squareness of the hysteresis loop is not good, this magnet is disadvantageous both in the temperature characteristics and in use.
Method (3) is a unique method in that the hot pressing process is used in two stages. However, it cannot be denied that this method is very inefficient in actual mass production.
Furthermore, according to this method, the crystal grains are remarkably coarse at a relatively high temperature of, for example, more than 800°C, which lowers the coercive force iHc to such a great extent that a practical permanent magnet is not obtainable.
A further method (4) of manufacturing a permanent magnet is disclosed in French patent publication no. 2 586 323. In this disclosure, a method is used, comprising melting an alloy, pouring the molten alloy into a mould casting it and either hot working before casting or applying a resin thereafter. This method results in a permanent magnet which overcomes the problems of oxidation and the need to handle a green body. The resultant permanent magnet also has slightly improved magnetic characteristics over method (1), (2) and (3) above. However, the improvement is not significant and given the importance of such magnets greater magnetic characteristics are considered essential. Moreover, such a magnet still suffers from the disadvantage that the coercive force is unacceptably low since the crystal grains are relatively tiny.
The present invention seeks to eliminate the above described disadvantages of known methods and to provide a magnetic alloy and a method of making the same which can readily be used to manufacture a high performance low cost permanent magnet.
According to one aspect of the present invention there is provided a rare earth-iron magnetic alloy as recited in claim 1.
According to a further aspect of the present invention there is provided a method of making a rare earth-iron magnetic alloy as recited in claim 11.
As described above, the known methods of making rare earth element-iron magnets have serious defects such as difficulty in handling pulverised powder and poor productivity.
To ameliorate these defects, methods of making magnets in a bulk state have been studied, and it has been found that it is possible to obtain a sufficient coercive force by casting a raw material containing a rare earth element, a transition metal and boron as the main ingredients so as to have a fine columnar macrostructure. To make a magnet, the cast material is hot worked so as to make the magnet anisotropic, and the magnet is then heat treated. To state this more specifically:

(1) since the macrostructure of the cast body has a fine columnar structure, it is possible to manufacture a magnet having plane anisotropy (the degree of alignment of the easy magnetisation axis is about 70%) merely by heat treating the magnet in the cast state;
(2) since the macrostructure of the cast body has a fine columnar structure, hot working accelerates the action of making the magnet uniaxially anisotropic, thereby considerably enhancing the degree of alignment of the easy magnetisation axis; and
(3) as a result of (1) and (2), it is possible to make a high performance magnet without the step of making the alloy into a powdered form, which is difficult to handle, so that heat treatment and strict atmosphere control can be omitted thereby enhancing the productivity of a magnet and greatly reducing the equipment cost.

The optimum composition of an R-Fe-B magnet is conventionally considered to be R₁₅Fe₇₇B₈, as is described in Reference 2 above. In this composition R and B are richer than in the composition R₂Fe_82.4B_5.9, which is obtained by calculating the main phase R₂Fe₁₄B, compound in terms of percentage. This is because a non-magnetic phase such as an R-rich phase and B-rich phase as well as the main phase is necessary in order to obtain a coercive force.
In contrast, in the appropriate composition in the present invention, the maximum value of the coercive force is obtained when the B content is lower than that of the main phase. Conventionally, this composition range has not been taken into much consideration, because the coercive force is greatly reduced when a sintering process is used.
However, when a casting process is adopted, the coercive force is easy to obtain when the B content is lower than the stoichiometric composition, and difficult to obtain when the B content is higher.
The reason for this phenomenon is thought to be as follows:
Whether the sintering process or the casting process is adopted, the coercive force mechanism itself conforms to the nucleation model. This is obvious from the fact that the initial magnetisation curves of the coercive forces in both cases show a steep rise such as those of SmCo₅.
The coercive force of a magnet of this type fundamentally conforms to a single magnetic domain model. In other words, in this case, if an R₂Fe₁₄B compound having a large crystal magnetic anisotropy is too large, the magnet has a magnetic domain wall in the crystal grains, so that reverse magnetisation is easily caused by the movement of the magnetic wall, thereby reducing the coercive force.
On the other hand, if the crystal grains are reduced in size to less than a certain dimension, no magnetic wall exists in the grains and since reverse magnetisation is caused only by rotation, the coercive force becomes large.
In short, in order to obtain an appropriate coercive force, it is necessary that the R₂Fe₁₄B phase has an appropriate grain diameter. The appropriate grain diameter is about 10µm (microns) and, in the case of a sintering type magnet, it is possible to determine the grain diameter by adjusting the grain size of the powder before sintering.
On the other hand, in the case of adopting a casting process, since the size of the crystal grains of the R₂Fe₁₄B phase is determined in the step of solidifying the molten material, attention must be paid to the composition and the solidification process.
Especially, the composition has a great influence on the grain size, and if not less than 8 atm % of B is included, the R₂Fe₁₄B phase as cast is apt to have coarse grains, so that it is difficult to obtain a good coercive force unless the quenching rate is increased more than usual.
On the other hand, in a low boron region, it is easy to obtain finer crystal grains by selecting an appropriate mould, controlling the casting temperature, etc. This region can be said from another point of view to be a phase richer in Fe than R₂Fe₁₄B phase, and Fe is first crystallised out as a primary crystal in the solidification step and subsequently the R₂Fe₁₄B phase appears by peritectic reaction. At this time, since the quenching rate is much higher than the equilibrium reaction, the R₂Fe₁₄B phase solidifies around the primary crystal Fe. In this composition region, since B is less, a phase richer in B such as an R₁₅Fe₇₇B₈ phase, which is a typical composition of a sintering type magnet, is almost negligible. The heat treatment is carried out in order to diffuse the primary crystal Fe so as to attain the equilibrium state, the coercive force largely depending on the diffusion of the Fe phase.
The reason why a columnar macrostructure is used in the present invention will now be explained.
As described above, adoption of the columnar structure has two effects; one is that it enables the permanent magnet to possess plane anisotropy, and the other is that it enables the permanent magnet to obtain a high performance during hot working.
The former effect will first be explained. The intermetallic compound R₂Fe₁₄B, which becomes the source of the magnetism of the R-Fe-B magnet, has the property of distributing the easy magnetisation axis C in a plane perpendicular to the columnar crystals when the columnar structures are grown. In other words, the C axis is not in the direction of columnar crystal growth but is in a plane perpendicular thereto, namely, the permanent magnet has anisotropy in a plane. This permanent magnet naturally and very advantageously has a higher performance than a permanent magnet which has a uniaxial macrostructure. Even if the columnar structure is adopted, the grain diameter must be fine in terms of the coercive force, and it is therefore desirable that the B content is low.
The adoption of a columnar structure further enhances the effect of hot working on bringing about anisotropy. If the degree of magnetic alignment is defined as:

$M.A. = B X / √ \bar{(Bx² + By² + Bz²)} x 100 (%)$

wherein Bx, By, Bz represent residual magnetic flux density in the directions x, y and z respectively, the degree of magnetic alignment in the isotropic magnet is about 60%, and in the plane anisotropic magnet, it is about 70%. The effect of hot working on bringing about anisotropy (enhancing the degree of magnetic alignment) exists irrespective of the degree of magnetic alignment of the material being processed, but the higher the degree of magnetic alignment of the original material, the higher the degree of magnetic alignment of the final processed material. Therefore, enhancing the degree of magnetic alignment of the original material by adopting a columnar structure is effective for finally obtaining a high performance anisotropic magnet.
A preferable composition range of a magnetic alloy according to the present invention will now be explained.
As a rare earth element, at least one selected from the group consisting of Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu is used. Pr produces the highest magnetic performance.
Practically, Pr, Pr-Nd alloy, Ce-Pr-Nd alloy, etc. are used. Addition of a small amount of an element, e.g. heavy rare earth elements such as Dy and Tb, and Al, Mo, Si, etc. enhances the coercive force.
The main phase of an R-Fe-B magnet is R₂Fe₁₄B. Therefore, if the R content is less than about 8 atm %, it is impossible to form the above described compound and the magnet has a cubic structure the same as that of an alpha-iron magnet, so that it is impossible to obtain high magnetic properties.
On the other hand, if the R content exceeds 30 atm %, a non-magnetic R-rich phase increases, thereby greatly lowering the magnetic properties. The appropriate range of R is therefore 8 to 30 atm %. In the present invention, since the magnet is made by casting, the R content is preferably 8 to 25 atm %.
B is essential for forming an R₂Fe₁₄B phase, and if the B content is less than 2 atm %, a rhomobohedral R-Fe structure is formed, so that a high coercive force is not expected. On the other hand, if the B content exceeds 28 atm %, a non-magnetic B-rich phase increases, thereby greatly lowering the residual magnetic flux density. As a cast magnet, the B content is preferably 2 to 8 atm %. If it exceeds 8 atm %, it is difficult to obtain a fine R₂Fe₁₄B phase, so that the coercive force is reduced.
Co is an effective element for increasing the Curie point of an R-Fe-B magnet. Basically, the site of Fe is substituted by Co to form R₂Co₁₄B, but this compound has a small crystal magnetic anisotropy and with increase in amount, the coercive force of the magnet as a whole decreases. Therefore, in order to provide a coercive force of not smaller than 1 KOe, to use not more than 50 atm % Co is preferable.
Al has an effect of increasing the coercive force (Reference 6: Zhang Maocai et al: Proceedings of the 8th International Workshop of Rare-Earth Magnets, p. 541 (1985)).
Reference 6 shows the effect of Al on a sintered magnet, but the same effect is produced on a cast magnet. However, since Al is a non-magnetic element, if the amount of Al to be added is increased, the residual magnetic flux density is lowered. If the amount exceeds 15 atm %, the residual magnetic flux density is lowered to not more than that of a hard ferrite and the roll of a rare earth magnet which has high performance is not attained. Therefore, the amount of Al to be added is not more than 15 atm %.
The invention is illustrated, merely by way of example, in the accompanying drawing, which is a flow chart of manufacture of an R-Fe-B magnet using a magnetic alloy according to the present invention.
An alloy having the composition shown in Table 1 was first melted in an induction furnace, and cast into an iron mould to form a columnar structure.
The casting was annealed at 1000°C for 24 hours to be magnetically hardened.
In the case of a casting type magnet, the casting was cut and ground in this stage, thereby obtaining a magnet having plane anisotropy obtained by utilising the anisotropy of the columnar crystals.
In the case of an isotropic type magnet, the cast body was subjected to hot working prior to annealing.
In this example a hot pressing method was used as hot working. The processing temperature was 1000°C.
The magnetic properties of each of the magnets obtained both by casting the respective alloys and by further hot pressing the cast alloys are shown in Table 2.

Both Pr₁₄Fe₈₂B₄₁ which had the highest performance among these and Nd₁₅Fe₇₇B_8, which was the optimum composition in the Reference 2, were cast into an iron mould to form a columnar structure, a vibrating mould to form an equiaxial structure and a ceramic mould to form coarse grains. The magnetic properties of the respective magnets were compared with each other, and the results are shown in Table 3.

TABLE 1

No.	Composition
1	Pr₈Fe₈₈B₄
2	Pr₁₄Fe₈₂B₄
3	Pr₂₀Fe₇₆B₄
4	Pr₂₅Fe₇₁B₄
5	Pr₁₄Fe₈₄B₂
6	Pr₁₄Fe₈₀B₆
7	Pr₁₄Fe₇₆B₈
8	Pr₁₄Fe₇₂Co₁₀B₄
9	Pr₁₄Fe₅₇Co₂₅B₄
10	Pr₁₄Fe₄₂Co₄₀B₄
11	Pr₁₃Dy₂Fe₈₁B₄
12	Pr₁₄Fe₈₀B₄Si₂
13	Pr₁₄Fe₇₈Al₄B₄
14	Pr₁₄Fe₇₈Mo₄B₄
15	Nd₁₄Fe₈₂B₄
16	Ce₃Nd₃P₈Fe₈₂B₄
17	Nd₁₄Fe₇₆Al₄B₄
18	Pr₁₄Fe₈₀Ni₂B₄
19	Pr₁₄Fe₈₀Cu₂B₄

TABLE 2

No.	Casting Type		Hot Working Type
	iHc(KOe)	(BH)max(MGOe)	iHc(KOe)	(BH)max(MGOe)
1	3.5	1.9	6.2	7.5
2	11.0	7.3	18.3	36.9
3	8.2	5.7	14.5	28.3
4	7.0	4.2	13.7	19.4
5	3.4	2.5	7.2	13.5
6	6.7	6.8	12.4	28.4
7	1.5	1.5	3.5	7.0
8	9.5	7.0	14.9	29.7
9	6.0	4.5	9.2	19.9
10	3.5	4.3	6.2	7.6
11	12.9	8.0	21.0	22.7
12	10.7	6.5	18.9	26.8
13	11.7	7.9	19.6	29.4
14	11.8	7.4	18.6	27.6
15	7.7	6.3	14.3	23.0
16	8.2	6.8	15.8	24.3
17	11.7	7.8	16.0	27.0
18	12.5	7.8	19.7	37.2
19	13.0	7.9	20.5	36.7

It is clear from Table 3 that the compositions containing a smaller amount of B show a higher magnetic performance.
In addition it is observed that all the magnetic properties such as the coercive force iHc, the maximum energy product (BH) max and the degree of magnetic alignment with a columnar structure formed by a method according to the present invention are more excellent than those of the other magnets both in the form of a casting and after being subjected to hot working.
As described above, according to a method of making a permanent magnet using a magnetic alloy according to the present invention, it is possible to obtain a high performance magnet merely by heat treating a cast ingot without pulverising it, thereby advantageously enhancing productivity.

Claims

A rare earth-iron magnetic alloy comprising an alloy of between about 8 to 30 atomic percent of at least one rare earth element, between about 2 and 8 atomic percent boron and the balance iron, wherein said magnetic alloy has a columnar macrostructure and an easy axis of magnetisation aligned in a plane perpendicular to the direction of crystal growth.
The rare earth-iron magnetic alloy of claim 1, wherein the iron is present in an amount up to about 88 atomic percent.
The rare earth-iron magnetic alloy of claim 1 or claim 2, wherein the rare earth element is selected from the group consisting of yttrium, lanthanum, cerium, praseodymium, neodymium, promethium, samarium, europium, gadolinium, terbium, dysprosium, holmium, erbium, thulium, ytterbium, lutetium and mixtures thereof.
The rare earth-iron magnetic alloy of any preceding claim, further including an effective amount of cobalt for increasing the Curie temperature of a magnet formed from the magnetic alloy.
The rare earth-iron magnetic alloy of claim 4, wherein the cobalt is present in an amount up to about 50 atomic percent.
The rare earth-iron magnetic alloy of any preceding claim, further including an effective amount of at least one coercive force enhancing member selected from the group consisting of aluminium, molybdenum, copper and mixtures thereof for enhancing the coercive force of a magnet formed from the magnetic alloy.
The rare earth-iron magnetic alloy of claim 6, wherein the coercive force enhancing member is present in an amount up to about 15 atomic percent.
The rare earth-iron magnetic alloy of claim 3, wherein the rare earth element is selected from the group consisting of neodymium, praseodymium, cerium and mixtures thereof, further including an effective amount of cobalt for increasing the Curie temperature of the magnetic alloy and an effective amount of at least one selected from the group consisting of aluminium, molybdenum, copper and mixtures thereof for enhancing the coercive force of a magnet formed from the magnetic alloy.
A rare earth-iron magnetic alloy composition comprising:
at least one rare earth element in an amount between about 8 and 30 atomic percent;
boron in an amount between about 2 and 8 atomic percent;
an effective amount of cobalt for increasing the Curie temperature of a magnet formed from the magnetic alloy;
an effective amount of at least one coercive force enhancing member selected from a group consisting of aluminium, molybdenum, copper and mixtures thereof for enhancing the coercive force of a magnet formed from the magnetic alloy;
the balance of iron; and
wherein the magnetic alloy is anisotropic and has a columnar macrostructure.
The rare earth-iron magnetic alloy of claim 11, wherein the rare earth element is selected from the group consisting of neodymium, praseodymium, cerium, and mixtures thereof, cobalt is present in an amount up to about 50 atomic percent and wherein the coercive force enhancing member is aluminium in an amount up to about 15 atomic percent.
A method of manufacturing a rare earth-iron magnetic alloy comprising:
casting a molten alloy composition including at least one rare earth element, iron and boron to form an anisotropic cast ingot having a columnar macrostructure.