IE871643L - Magnet manufacture - Google Patents

Description

The present invention relates to a method of making a permanent jsagnet from an alloy comprising a. rare earth, iron and boron (H-Fe-B)- The term "rare earth™, as used in this specification is to be understood to include yttrium (¥).

At present, the following three methods are in practice used for manufacturing a magnet from such an alloys-CD A sintering method based upon a powder metallurgy technique (Reference No. 1) (2) A method in which rapidly quenched ribbon fragments having a thickness of about 30 microns are prepared by a melt spinning apparatus and are used for producing an amorphous alloy„ the magnet being made from the ribbon fragments by a resin-bonding technique (Reference No. 2); and (3) A method in which a mechanical alignment treatment is performed on the fragments prepared by the above method (2) by a 2-step hot pressing technique (Reference No. 2).

Reference No. 1: M. Sagawa, S. Fujimura, N. Togava, H. Yamamoto and Y. Matsuura; Japan Applied Physics Vol. 55(6), 15 Harch 1984, p2083.

Reference No. 2; R.W. Lee; Applied Physics Letters, Vol 46 (8), 15 April 1985, p790. The prior art techniques mentioned above are explained below.

In the sintering method (l)f an alloy ingot is made by melting and casting, and then the ingot is pulverised to a fine powder whose particle diameter is about 3 saicrons. The powder is then kneaded with a binder -which serves as a Moulding additive and is press-moulded in a magnetic field to obtain a moulded body. The moulded body is sintered in an argon atsaosphere for 1 hour at about 1100°C and thereafter quenched to room temperature. Mtar sintering, the body is heat-treated at about S00°C, thereby increasing its intrinsic coercivity.

In the method (2), quenched ribbon fragments of R-Fe-B alloy are prepared by a sielt spinning apparatus -shich spins at an optimum substrate velocity. The obtained fragments are ribbon-shaped, and have a thickness of 30 microns, and consist of an aggregation of grains whose diameter is 1000 Angstroms or less. These fragments are fragile and are saagnetically isotropic since the grains are distributed isotropically. The fragments are crushed to form particles of suitable size and these particles are kneaded with resin and are press-moulded. At this tijse, the density of the material is 85 volume percent under a pressure of about 2 7 tons/cm » In the method (3) , the rapidly quenched ribbons or . fragments are passed into a graphite or other suitable high-temperature die which has beers pre-heated in a vacuum or in an inert gas atmosphere to about 700°C. When the tasaperetura of the ribbons rises to the predetermined temperature» the ribbons are subjected to a unidirectional pressure. The temperature and the tiae are not limited although it is desirable that the temperature should be 725 ± 250°C 2 and that the pressure should be F ~ 1.4ton/esi so as to achieve sufficient plasticity. At this ti»se, the grains of the magnet are slightly aligned in the pressing direction but are isotropic as a -whole.

A further hot pressing process is performed using a larger cross-section die. As a rule„ the hot pressing is effected at a temperature of 700°C and at a pressure of o 0.7 ton/cm" for a period of several seconds. Then the thickness of the material is reduced by half of the initial thickness and a xsagnetic alignment is introduced parallel to the direction of pressing so that the alloy becomes anisotropic.

The above process is called a ^two-stage hot-press procedure". By using this procedure an anisotropic R-Fe-B series magnet of high density is produced.

It is preferable to sake the particle diameter of the grains of the ribbon fragments initially prepared by the xa-slt spinning saathod slightly smaller than the grain diameter at 'which the taaximuin intrinsic coercivity is produced- This is because„ since the grains become coarse to some degree during the hot-press procedure,, if the grain diameter of the crystals before the hot-press procedure are & little smaller than the optimum diameter, they will be the optimum sise after the procedure.

The prior art techniques mentioned above, however, have some disadvantages., In the sintering method (1), the alloy must be ground to a fine powder. However,a R-Fa-3 series alloy is extremely easily oxidized and so a R-Fe-B series alloy powder is all the nsore easily oxidised. Accordingly, the oxygen concentration of the sintered body is inevitably high.

Furthermore, when moulding the powder, an additive such as, for example, zinc s tear at a is necessary. Although such an additive is eliminated prior to the sintering process, soxae of the additive remains in the final magnet in the fone of carbon* Such carbon greatly deteriorates the magnetic performance of a R-Fe-B magnet.

The moulded body, after the moulding with the moulding additive has been effected, is called the "green body" and this is easily breakable and very difficult to handle.

Therefore, it is very difficult to place the green bodies in the sintering furnace in good order, which is a big disadvantage.

Because of the above disadvantages, in order to manufacture R-Fe-B series sintered magnets„ expensive equipment has been necessary. Moreover, productivity has been extremely low, which has resulted in high manufacturing costs for this type of magnet.

Thuse the sintering method (1) is not a satisfactory method for making good 'use of the advantage of the cheap raw material cost of an R-Fa-B series magnet.

In methods (2) and (3), a vacuum melt spinning apparatus is used. At present,, however, this apparatus has a low output and is expensive.

In method (2), the crystals of the resulting magnet are isotropic and so the energy production is low and the squareness of the hysteresis loop thereof is not good. Therefore the s&gnet produced by the method (2) has bad temperature coefficients and is unsatisfactory in practical use.

The method (3) is performed in two stages and is therefore very inefficient.

According, therefore, to the present invention, there is provided a method of making a permanent magnet comprising melting and casting an alloy whose constituents comprise at least one rare earth, iron and boron; hot working the cast alloy ingot at a hot working temperature of at least 500°C and a a strain race de/dt in the range of 10~4 to 1 per second, where £ is the logarithmic strain and t is time; and forming the hot worked alloy ingot into a permanent magnet.

The logarithmic strain £= £ . £«/£,, where £, is the n i l I length of the cast alloy prior to the said hot working, and 12 is its length after the said hot working, being the natural logarithm £oge.

The permanent asacnet is preferably annealed at a temperature of at least 2S0°C.

The hot worked alloy may be pulverised to forsa a powder, 7 the powder being kneaded with an organic binder.

The said hot working temperature is preferably in the range of 800°C to 1050°C.

The hot working preferably effects a reduction ratio 5 of at least 60% in the thickness., cross-sectional area or diameter of the cast alloy ingot. Thus if the reduction is a reduction in the thickness, the reduction ratio may be represented as d^ - d^ x 100 (%), where d^ is the thickness dl of the cast alloy prior to the said hot working and d^ is its thickness after the said hot working.

The alloy may contain a total of 8 to 30 atomic percent of one or more rare earths.

The alloy may contain 2 to 28 atomic percent of boron. 15 Thus the alloy may, for examples contain 8 to 30 atomic percent of Rs 2 to 28 atomic percent of B, 50 or less atomic percent of Cos 15 or less atomic percent of A I, the balance being iron and impurities which are inevitably included during the preparation process. 20 The invention also comprises a method of making a permanent magnet comprising melting and casting a Ce^Pr^Ndj^Fe^^^ alloy; hot working the cast alloy ingot at a hot 'working temperature of at least 500°C„ and a -4 strain rate de/dt in the range of 10 to 10 per second where e is the logarithmic strain and t is time; and forming the hot worked alloy ingot into a permanent magnet. 8 The present invention further comprises a method of Baking a permanent magnet comprising malting and casting an alloy whose constituents comprise sc lease one rare earth, iron and boron; hot working the cast alloy ingot at a hot working temperature of 1000°CS and at a strain rate d £ /dt in ~4 the range of 10 to 10 per second where £ is the logarithmic strain and t is time; and forming the hot worked alloy ingot into a permanent magnet, As described before, the known methods for preparing a 10 rare earth-iron series permanent magnet, i.e. the sintering method and the quenching method, have the respective disadvantages that handling of the powder is difficult and that productivity is poor.

It has been foundf however, that if the alloy has the preferred composition referred to in the previous paragraph, the alloy will be faade fine and anistropic by the hot working. The hot working on the cast .ingot at the temperature of 500°C or higher may be such as to produce fine crystal grains whose grain axis is aligned in a specific direction, thereby soaking the cast alloy -magnetically anisotropic.

In order to improve the magnetic properties and especially to increase the intrinsic coercivity of the resulting magnet, the alloy may consist of 8 to 25 atomic percent of R, 2 to 8 atomic percent of B, 40 or less atomic percent of Co, 15 or less atomic percent of Ai, the balance being iron and impurities which are inevitably included during the preparation process, the resulting magnet being magnetically hardened by heat-treatment at a temperature of 250°C or mora. In this case, sufficient intrinsic coercivity is obtained merely by the hot working.

For resin-bonded magnets, an alloy of the above composition may be pulverised to a fine powder by utilising its property of easily producing a hydrogenated compound, the said fine powder being kneaded with an organic binder and then cured to obtain a resin-bonded isagnet.

To obtain a resin-bonded magnet. by the usual pulverization, the povder raav be such that, by utilising the property that the grains are easily made fine by the hot working, each grain of the powder includes a plurality of magnetic R Fe B grains even after the pulverization has been effected, and the powder is kneaded with the organic binder and cured to obtain the resin-bonded magnet.

In the method of the present invention, the hot working, which isay be effected to sake the ingot anistropic, stay be only a one-stage process and not a two-stage process as in the quenching method disclosed in Reference No. 2, Moreover, the intrinsic coercivity of the product will be substantially increased because the grains are made fine. Furthermore, since it is not essential to pulverize the cast ingot, it is not necessary to strictly control the atmosphere for sintering and so on, thereby greatly reducing the equipment cost.

It is another advantage of the present invention that the resin-bonded magnet which can be obtained by the method of the present invention is not originally isotropic as is the jaagnet obtained by the usual quenching method and that an anisotropic resin-bonded magnet can easily be obtained.

Thus the advantages of B-Pe-B magnet of high-performance and low cost can be achieved.

A report on the magnetisation of an alloy in the bulk state has been presented by Hiroaki Miho et al. (The Lecture Meeting of The Japanese Institute of Metals „ Autumn 1985, Lecture Mo. 544), However, this report refers to ssaall alloy sasples having a composition of NdJg 2Fa5o 7C02? 6V* 3 B9 which are melted in air, exposed to an ergon gas spray and then extracted for sampling. Accordingly, it is considered that, in the study of this report,, the fine grains obtained by quenching occurred because of the small-quantity sampling.

Me have found, however, that, in the composition disclosed in this report, the grains of the main phase Nd^Fa^B thereof become coarse when they are cast by the usual casting method, Although it is possible to make an alloy of the composition Nd 2Fe50 7C°22 5V' 339 2 anisotroPic hot working, it is very difficult to obtain sufficient intrinsic coercivity as a persaanent magnet for the resulting body.

Also, ve have found that, in order to obtain a magnet of sufficient intrinsic coercivity even by the usual casting method, the composition of the starting material should have a B-poor composition of 8 to 25 atomic percent of E, 2 to 8 atomic percent of 3, 50 or less atomic percent of Co, 15 or less atomic percent of hX, the balance being Fe and impurities.

A typical optimum composition for a R-Fe-B series magnet of the prior art is believed to be R^Fe^Bg as shown in Reference No. 1. In this composition, R and B are richer than in the composition R, ™Fe00 0 'which is equivalent X I • / o/. • 4 D <->" in atomic percentage to the main phase R^Fe^B coajpound.

This is because, in order to obtain sufficient intrinsic coercivity, not only the sain phase but also the non-magnetic phase of R-rich phase and B-rich phase are necessary. 1 2 In the composition B-poor region referred to in the preceding paragraph but one, the intrinsic coercivity becomes a maxiimusi when B is poorer than in the usual composition. Generallyt such a 3-poor composition shows a great decrease in intrinsic coercivity when the sintering method is applied and therefore this composition region has not been considered very carefully in the past. However„ if the usual casting method is employed,an high intrinsic coercivity is obtained only in the said composition region and, in the B-rich composition which is the main composition region for the sintering method, the intrinsic coercivity is not sufficient.

The reason for the above is considered to be as follows. Primarily, by employing either a sintering method or the casting method of the present invention, the intrinsic coercivity mechanism of the magnet itself is in accordance with a nucleation model- This is proved by the fact that the initial magnetisation curve of the aagnets by both methods shows a steep rise similar for example, to that of SsCOj. „ K magnet of this type has the intrinsic coercivity according to t^s single domain saodel. Namely, if the grains of E^Fe^B compound having a large crystal magnetic anisotropy are too large*, magnetic domain walls are introduced into the grains and accordingly movement of the magnetic doscain walls causes the reverse magnetisation 1 3 to be easily inverted,, thereby decreasing the intrinsic coercivity. Hhereas, if the grains of R,Fe1^B compound are smaller than a specific sis a, the snagnetic walls disappear from the grains® In this esse, since the reversal of the saagnetiaation is caused only by the rotation of the isagnetisation, the intrinsic coercivity decreases.

Thus, in order to obtain sufficient coercivity, the 1 phase is required to have adequate grain diameter, i.e. about 10 microns. When the sintering method is applied, the grain diameter can be suitably adjusted by the adjustment cf the powder diameter before sintering. However, by the casting aethod, the grains diameter of the ^2Fei4B compound is determined ■when the liquid snaterial is solidified. Accordingly, it is necessary to control the composition and the solidification process with great care.

The composition is especially important. If the B content of the alloy is more than 8 atomic percent, it is very likely that the grains of the R^Fe^B phase in the magnet after casting are larger than 100 microns. Accordingly, in this case, sufficient intrinsic coercivity is difficult to obtain in the cast state without using a quenching device as used in Reference 2. On the contrary,, in the said B-poor composition region, the diameter of the grains of the signet is easily reduced by adjusting the kind of mould, the moulding temperature and so on. However, in either case, the I 9. grains cf the xnain phase J^Fa^B are made finer by performing the hot working and so the intrinsic coercivity of the magnet increases after the hot working.

The composition' region in which there is sufficient 5 intrinsic coercivity in the cast state, i.e. the said 3-poor composition, can also be called a Fe-rich composition« At the solidifying state, Fe first appears as the primary phase and then the R^Fe^B phase appears by peritectic reaction. At this time, since the ccoling speed is much higher than the 10 speed of equilibrium reaction, the sasple is solidified in such a way that the R^Fe^B phase surrounds the primary phase Fe. Since this composition region is 3-poorf the B-rich phase as seen in a R^Fe^Bg magnet which is the typical composition suitable for use in the sintering method is necessarily so 15 srca.ll in quantity that the B-rich phase can almost be neglected.

The heat treatment employed for the said B-poor composition is used to diffuse the primary phase Fe and to attain an equilibrium state, so that the intrinsic coercivity of the resulting magnet greatly depends on the diffusion of Fe. 2Q A resin-bonded magnet is actually prepared by the quenching method of Reference No. 2. However, since the powder obtained by the quenching method consists of isotropic aggregation of polycrystals whose diameter is 1000 Angstroms or less f the powder is magnetically isotropic. Thus, an anisotropic magnet 25 cannot be obtained and the advantages of a R-Fe-B series aiagnetf i.e. low cost and high performance, are not achieved according to the quenching method. When a R-Fe-B series magnet is to be prepared, the intrinsic coercivity of the magnet is kept sufficiently high by a pulverizing step comprising hydrogen decrepitation which causes little mechanical distortion, and accordingly the resin-bonding can be achieved. The greatest merit of this method is that an anisotropic magnet can be prepared otherwise than by the method of Reference No,2.

There are two reasons why a resin-bonded R-Fe-B series magnet can be prepared by pulverising the alloy by using its property of easily producing an hydrogenated compound, kneading the resulting powder with an organic binder, and curing the alloy to obtain a resin-bonded asagnet.

First, attention should be given to the fact that the critical radius of the single fiomain of the 4® compound is ssuch smaller than that of SraCo„ and so on and is of the order of submicrons. It is extremely difficult to pulverise material to such a small grain diaaaeter by the usual mechanical pulverisation. Moreover, the obtained powder is too highly activated and consequently is oxidised and ignited very easily, and therefore the intrinsic coercivity of the resulting nsagnet is very low for its grain diasieter. hs a result of a study of the relationship between the grain diameter and the resulting intrinsic coercivity, however, we have found that the intrinsic coercivity was at most, sosie kOe and did not increase even by performing surface treatment on the magnet.

Another problem is a distortion caused by mechanical working. For example, if a magnet having an intrinsic coercivity of iOkOe in the sintered state is mechanically pulverized, the resulting powder of a grain diameter of 20 to 30 microns possesses a coercivity as low as less than 1 kOe. If a SmCo magnet, which is considered to have a similar nucleation model, is mechanically pulverized, such a decrease of the intrinsic coercivity dees not occur and a powder having sufficient coercivity is easily prepared. The reason for such a phenomenon is assumed to be that the effect of the distortion and so on caused by pulverisation and working on a H-Fe-B series magnet is considerable. This effect is a critical problem when a small magnet such as a rotor magnet of a step motor for a watch, is cut from a sintered magnet block.

For the above reasons, namely, that the critical radius is small and that the effect of mechanical distortion is considerable, a resin-bonded magnet cannot be obtained by the usual pulverization. In order to obtain a powder having sufficient intrinsic coercivity, a powder whose grains include a proportion of H_Fe B grains as disclosed in Reference J«o. 2 should be prepared. However, the quenching method of Reference Ho. 2 has low productivity, Furthermore,, it is actually impossible to prepare a powder of this kind by pulverisation of a sintered body because the grains grow larger to sosae degree during sintering and it is necessary to sake the grain diasaetar before sintering smaller than the diameter finally desired. However, if the grain diameter of the powder is so small, the oxygen concentration thereof is esctremelv high and the performance of the aiagnet is far from satisfactory.

Thus e at present, the allowable grain diameter of the K„Fe„ .B compound after sintering is about 10 microns, but 2 14 the intrinsic coercivity is reduced to almost sero after pulverisation.

We have also considered the ssatter of the fining of grains by hot working. It is relatively easy to make a R^Fe^B cosipound in the aoulded state of about the same size as the grains prepared by sintering. So, by performing hot working on a cast block having a RFeB phase of that grain size, the grains are made finer and aligned, and are pulverized thereafter. By such a method, since the grain diameter of a powder for a resin-bonded aiagnet is between 20 and 30 microns, it is possible to include a plurality of R-Fe,„B grains 4 JL*5 in the powder f so that a powder having sufficient intrinsic coercivity is obtained. Moreover, these obtained powders are not isotropic as they are when obtained by the quenching method of Reference Mo. 2 but can be aligned in the laagnetie field, and accordingly* an anisotropic magnet can be prepared from the powder of this type. Of course, if the grains are 1 8 pulverised by hydrogen decrepitation, the intrinsic coercivity is maintained better.

The reason for the particular composition of the alloy used in the present invention is explained below.

As the rare earth elements used in the said alloy, one may employ one of iaore of the elements Y, La, Ce, Pr, Ho, Nd, Sm, Eu, Gd, Tb, Dy, Er, Tm, Yh and Lu, The highest i&agnetic performance is obtained when Pr is selected either alone or in combination vith other rare earth elements. Accordingly, for the practical use,, a Pr alloy, a Pr-Nd alloy,, a Ce-Pr-Nd alloy and so on are used.

A small amount of an additive consisting of heavy rare earth elements such as Dy, Tb etc., say be emploved in the a!loy. Moreover, RS., Mo, Si and so on are sometimes desired in the alloy in order to enhance the intrinsic coercivity.

The main phase of the R-Fe-B series magnet is R2Fe14B. If R is less than 8 atomic percent, the above compound is not present but a body centered cubic compound of the same structure as a-iron is present and consequently high magnetic properties are not obtained. If, on the contraryf, R is sxsore than 30 atomic percent? the quantity of the non-magnetic B-rich phase increases and the magnetic properties are extremely reduced. Thus, the amount of R is preferably between 8 and 30 atomic percent. However„ the range of R which is aara satisfactory for a cast nsagnet is between 8 and 25 atomic percent. i 9 B is an essential element to produce the R Fe B <£. 1*8 phase. If B is less than 2 atomic percent, the rhoashohedral R-Fe series is present and so high intrinsic coercivity is not obtained. However, as in the case of a magnet produced by the sintering method of the prior art„ if B is included in an axoount in excess of 28 atomic percent, the non-magnetic B-rich phase increases and the residual magnetic flux density is substantially reduced. The preferred upper limit of the arsount of B for a moulded magnet is therefore 8 atomic percent. If B is snore than 8 atomic percent, the fine H^Fe^B phase is not obtained unless a specific cooling is performed, and the intrinsic coercivity is low.

Co is an element which is effective to enhance the Currie point and has the effect basically to substitute the site of the Fe element so as to produce H^CO^^B. However, this R^Co^B compound has a small crystalline anisotropv field, and the more the compound C^Co,, increases, the less is the intrinsic coercivity of the aiagnet. Accordingly„ in order to obtain a coercivity of at least IkOe, which is considered to be sufficient for a permanent aiagnet, Co should be 50 atomic percent or less.

Ai has the effect of increasing the intrinsic coercivity as described in Reference No 4; Shang Maocai et al. Proceedings of the 8th international Worship on Rare-Earth Magnets, 1985, F541. This reference Ko 4 refers only to the effect of A& in the case of a sintered usagnet, although the saoe effect is present in case of a cast magnet.

Ho waver, since A& is a non-magnetic element, if the amount of Afc is large, the residual magnetic flux density decreases and if snore than 15 atomic percent thereof is present, the residual magnetic flux density is reduced to the level of hard farrite. Such a magnet does not achieve the high performance of a rare earth magnet. Therefore, the amount of MJL is preferably 15 atozaic percent or less.

The invention is illustrated, merely by way of example, in the accompanying drawings, in which:- Figure 1 illustrates alternative steps in a method of producing a magnet in accordance with the present invention; Figure 2 illustrates the alignment of the grains of a magnetic alloy by extrusion hot working, 1 being a hydraulic press, 2 being a die, 3 being a magnetic alloy, 4 being an arrow showing pressure, and being arrows showing the direction of the easy magnetization of the saagnetie alloy; Figure 3 shows the alignment of the grains of the isagnetie alloy by rolling hot working, 6 being rollers, 7 being a magnetic alloy, 8 being arrets showing the direction of rotation of the rollers, 9 being an arrow showing the direction of the movement of the asagnetic alloy, and being arrows showing the directions of easy magnetisation; and Figure 4 shows the alignment of the grains of the aagnetic alloy by stamping hot working, 11 being a stamp, 12 being a magnetic alloy f 13 being a base plate, 14 being arrows showing the direction of easy magnetisation, being an arrow showing the vertical movement: of the stamp 11, and IS being an arrow showing the direction of movement of the base plats 13.

Sxaapla 1 Reference is made to Figure 1 which is a flow diagram illustrating a method of malting a permanent magnet in accordance with the present invention.

First, an alloy of the desired composition is soelted in an induction furnace and is cast into a die to form a cast alloy ingot. 'Then, in order to give anisotropy to the sua g net j, various kinds of hot workings are per formed on the ingot. In this example, use was issade not of a general moulding method, but of a specific moulding method, that is, the Liquid Dynamic Compaction method (Reference No. 5; T.S. Chin et al„ J. Appl. Phys, 59(4)s 15 February 1986, P. 1297) whieh has a substantial effect in producing fine crystal grains by quenching.

The hot working method used in this example is one of the following:- (1) extrusion-type (Figure 2), (2) rolling-type (Figure 3) and (3) stamping-tvpe (Figure 4) each being performed at a temperature of 1000°C. In each -3 -2 case, the strain rate is from 10 to 10 per second and the reduction ratio is 80%.

In the case of the extrusion-type, in order to apply an isostatic pressure to the seuaple, a means to apply pressure to the sasxple from the side of the die was also provided. In the case of the rolling-and the stamping-types , the speed of rolling or stamping was adjusted so as to minimize the strain rate. Whichever type thereof is used, the axis of easy magnetization of the grains is substantially parallel to the direction into which the alloy is urged.

Alloys Nos 1 to 14 of the compositions indicated in Table 1 were melted and made into magnets by the process shown in Figure 1. The hot working applied to each sample is shown in the Table.

Hie annealing after the hot working was performed at a temperature of 1000°C for 24 hours. 2 3 TABLE 1 Alloy NO Composition hot working 1 Kd8 Fe84 B8 extrusion 2 N<J15 Fe77 38 rolling 3 Nd22 FS70B8 stamping 4 SWd_,. Fe,__ B, _ extrusion 58 12 Ce3.4 Nd8.5 Pr7.1 Fe75 B6 rolling 6 Nd17 Fe60 C°17 B6 stamping 7 Nd17 Fe60 C°15 ^2 B6 extrusion 8 Ce. Nd_ Pr, Fecc Co,c a£c 3, 4 9 o 55 15 5 6 rolling 9 Ce3 Ndl0 Pr8 Fe52 Co15 H°4 B8 stamping Ce3 Nd10 Pr8 Fe52 Co17 m2 B8 extrusion «j e Ce3 Nd6 Pr10 Fe54 Co17 Ta2 B8 rolling 12 Ce3 HdS Pr8 Fe50 Co19 "Sx2 B12 stamping 13 Ce3 Ndl0 Pr6 Fe50 C°15 Zr2 B14 extrusion 14 Ce3 Hd10 Prfi Fe56 C015 Hf2 Bg rolling The properties of the resulting magnets are shown in Table 2. For the purposes of comparison, the residual aagnetic flux density of a sample "after casting",, i.e. a sample on which the hot working has not been performed, is also shewn. 2 4 TABLE 2 All efter hot working after .casting off 4 BrCKG) 9.3 9.7 8.5 8.4 .8 j i 3 11:9 11. S Ii.5 9.2 n 9.6 9.1 13 Ml 8. bHc(KGOe) i?) if* A. a <3 « *£ 9 9 (W? A 1*J JH e *S! « 0 3.8 3.9 .1 10.7 0 ^ A S/9 6.9 I.I 8.0 5.8 7.3 (BH)nax(HGOe) | BrCKG) o. s 4.9 e i ■3 <* (*§ C 1 3.1 .5.8 5.8 .1 %!.% 28.8 15.8 13.2 ll.S .1 0„8 1.3 1.7 n o <»L » £" 1.4 8.1 6.1 6.0 (C i-1 & * D •4 7 "2 « P 4:3 6.2 ri sax inuuc 0.1 0.3 0.5 0.2 0.3 0.4 2.3 ti 9 «> o <W «"> M 1.a 2.3 S.2 1.9 S.l From Table 2, it is obvious that irrespective of whether the hot working is by extrusion rolling or stamping, the residual Magnetic flux density increases, end thereby the samples are stade magnetically anisotropic.

Example 2 The present example eaplovs a general casting method- First alloys Nos 15 to 31 of the composition indicated in Table 3 were melted in an induction furnace and cast into a die to form columnar zones. Hot working was effected at 1000°C using hot pressing , the strain rate being maintained -3 -2 between 10 and 10 per second, the reduction ratio being 80%. An annealing treatment was performed on the ingot at 1000°C for 24 hours to magnetically harden the same.

After the annealing, the siean grain diameter of the sample was about 15 microns.

In contrast, in the case of a cast magnet of the desired shape which has been made without hot working, a plane anisotropic magnet utilizing the anisotropy of the columnar zone is obtained. In order to produce a resin-bonded magnet in a 18-8 stainless steel container at room temperaturee hydrogen absorption in an hydrogen atmosphere of about 10 atmospheric pressure and hydrogen desorption at a pressure of 10 ^ torr were repeated and the samples were pulverised, and 4 weight percent of epoxy resin was kneaded thereafter. Then the compacts were moulded in a magnetic field of lOkOe which was applied 2 6 erpendieularly to the pressing direction.

Ths properties of the resulting magnets are shown able 4.

TABLE 3 Alloy No. composition Pr FeD, 8© B4 IS prift J.O Fe 80 B4 17 Pr 22 * 74 B„ 18 Pr F«=> 70 B„ 19 Pr 13 Fe B2 1>r ~~13 Fe81 B, o 21 Pr13 Pp " 79 B8 22 Pri2 C°10 B4 23 Pr12 59 C°25 B4 24 Pr13 Fs43 C°40 B4 Pr 13 **3 Fa 80 B4 26 Pr Fa 78 B. Si 4 2 27 Pr , la F® 76 hl& 3 4 28 Pr 16 76 Mo„ 3 4 29 Kd14 Fa78 "D a " 4 a 4 Ce3 'M/j Po p 3 *~10 80 4 31 Hd12 F®80 ^4 B4 2 7 TABLE 4 Alloy No. hot working type 17 18 19 21 22 23 24 26 27 28 29 31 iHcO 6.1 H ^ s JL %) & i> ii k C S .2 i.I 12.0 8.1 18. i 7 R ti' « M 3.8 18.0 15.9 18 .4 16.6 9.6 11.6 IS.? ax CMGOe] resin-bonded type T~1 n * < Mukuo B a n^> if v If 9.8 4.9 i .3 27.1 12.0 17.8 17.5 9.4 9.8 11.0 8.2 . 5.S 3.0 3.0 1.8 21.5 9.0 14.2 2.2 4.7 11.8 . a .5 18.8 18.2 8.2 9.7 12.8 2.9 7.7 28.8 13.4 17.4 24.5 12.5 IS.2 .4 13.0 at A <•» 16.D .1 13.3 « BB «' !|i S ill & n ® <* 12.5 15.0 7.5 9.3 .3 13.5 21.1 13.5 14.S hs will be seen, (BH)iaax and iBc are greatly increased by the hot working. This is because the grains are aligned by 'the hot working and the squareness of the BH curve is very such improved. In the quenching method of Reference Sc. 2„ on the other hand, iHc tends to be lowered by the hot working. Accordingly, it is one of the major advantages of the present invention that the intrinsic coercivity is generally improved.

TABLE 5 Pv pa n 81 4 Ce3Pri0Nd pea 3 "73 4 (d£/dt/second) 1H P(KOe) Br (KG) .iHc(MOe) Br (KG) io"5 - "4/s = 6 .9 9.0 .7 i io"4 ~ ~3/s 14.3 .8 .4 .S ~3 ~ 2 ~/s .1 9,5 12.0 .S io~2 - " Vs 16.0 8.8 13.8 8 « 9 j io-1 ~ 1/s IS.6 7.0 .9 7.1 i 1 - I /s X X .9 6.8 o i 102/s X X X X 2 103/s X X X X : the sansple was cracked. 29 Table 5 above shows the relationship between the strain rate and the magnetic properties of two alloys whose composition is indicated in Table 5. These alloys„ which constitute representative examples, are melted in an induction furnace and are cast in a die made of iron, and then they are subjected to hot pressing at various speeds. The results are shown in Table 5. In this case, the temperature is controlled at 1000°C, and the ratio of reduction is 80%. After the hot pressing, an annealing treatment is performed at 1000°C for 24 hours.

As shown in Table 5, when the strain rate is smaller than IO"'* per second, the intrinsic coercivity is decreased extremely. The reason for this is considered to be that the growth of crystal grains is hastened by the heat and that they are too bulky. On the other hand, the strain rate is too small, the productivity is decreased and thus the manufacturing cost is increased. When the strain rate is bigger than 1 per second, some samples are cracked due to the composition thereof, and such samples cannot be manufactured. Therefore, the desirable strain rate is frora 10 ^ to S per second, and that from 10 to 10 per second is desirable in order to get excellent magnetic properties.

TABLE 6 composition property plastic processing temperature C® C) room temper^* ature. 250 500 700 800 900 950 1800 1050 1 1100 I 1150 Pri7Fe?rB4 iHc(KOe) A A IO. 8 11.8 J 8.4 8.0 8.8 7.8 6.2 2.1 ) 70 C axis orientation^ A A 82 80 85 85 95 96 97 95 M30F®55 BIS iTtfcC iHc(KOe) X X .0 13.6 12.8 12.8 12.0 12.4. 9.2 .8 1.5 . c axis orientation rat 2 X X n 71 82 98 88 93 97 91 11.4 73 2.5 K KVfI0 IHc(KOe) X /\ .0 19.0 14.4 18.8 14.2 11.2 ,8 ] Fe50Co17Zr2 Ba C axis■ori~ X . ...

A 63 81 89 SB . 95 91 95 69 X s the sample cannot be 'manufactured A i the sample is cracked and cannot be measured. 3 1 The alloys of the -three compositions indicated in Table 5 as representative examples, are melted in an induction furnace and ere cast in a die made of iron. Then hot working is effected by extrusion at the various temperatures shown, and an annealing treatment at 1000°C for 24 hours is performed,, The relationship of the manufacturing temperature, the intrinsic coercivity and the C axis orientation rate are shown in TaJble S. The strain rate is arranged *-3 -2 to be from 10 to 10 per second, and the reduction ratio is arranged to be 80%.

Further,the C axis orientation rate represents the rate (volume %) of the easy magnetization axis of the crystal grains (corresponding to the C axis of the permanent magnet of the present invention) which are aligned ixi the same direction. The bigger this rate is, the finer the anisotropic magnet which can be obtained.

As shown in Table 6, if the manufacturing temperature is less than 500°Cr the sample is cracked and thus cannot be manufactured. The C axis orientation rate of 80% is desirable to obtain excellent magnetic properties. In this case, the manufacturing temperature should be from 800°C to 1100°C. However, if the temperature is 1100°C, the intrinsic coercivity is substantially decreased,, Thereforec, the best manufacturing temperature is from 800°C to 1050°C.

TABLE 7 ratio of Pr,? Fe?g B4 P "55 15 reduction (%) c axis orientation C axis (KOe) rate (%) (KOe) orientation rate % 0 4.3 58 • .5 60 4.7 68 6.7 56 40 4.9 71 7.4 70 SO 6.9 80 9.0 81 70 7.7 90 .8 93 80 8-6 96 12.4 98 90 9.4 95 .8 98 The alloys of "the compositions indicated in Table 7 are melted and cast by the same saethod as those of Table 5. Than hot working is effected by an extrusion raethod in this case at a temperature of 1000°C, the reduction ratio being varied. Then an annealing treatment, at a temperature of 1000°C for 24 hours is performed. The relationship of the reduction ratio, the intrinsic coercivity and the C axis orientation rate axe shown in Table 7. Further, the strain rata Is arranged to be from „3 =2 to 10 per second. As shown in Table 7, if the C axis orientation rate is required to be 80% or higher, the reduction ratio should be 60% or higher.

The present invention thus enables a pensanant magnet of sufficient intrinsic coercivity to be provided raerely by carrying out a hot working without pulverising the ingot as in the usual sintering method.

Furthermore, the hot working is only a one-stage process and is not a two-stage process as in the quenching method. Moreover, the hot working not only makes the saagnet anisotropic but also increases the intrinsic coercivity.

Thus the present invention provides a jnethod of making a permanent magnet which is such simpler than the sintering method or the quenching method in the prior art- Moreover, by hydrogen decrepitation or pulverization of the samples after the hot working, an anisotropic resin-bonded aiagnet can also be provided in accordance with the present invention,

Claims (1)

1. CLAIMS 1. a method of raaking'a permanent: magnet comprising malting and casting en alloy vhosa constituents comprise at least one rare earth, iron and boron,- hot working the cast alloy ingot at a hot working temperature of at least 500°C , and at a strain rate de/dt in the range of 10 co 1 per second where £ is the logarithmic strain and t is time; and forming the hot worked alloy ingot into a permanent magnet, 2. A method as claimed in claim 1 in -which the permanent magnet is annealed at a temperature of at least 250°C. 3. A method as claimed in claim 1 or 2 in which the hot worked alloy ingot is pulverised to form a powder, the powder being kneaded with an organic binder, 4. A method as claimed in any preceding claim in which the said hot working temperature is in the range of 800°C to 1050°C. 5. A method as claimed in any preceding claim in which the hot working effects a reduction ratio of at least 60% in the thickness, cross-sectional area or diameter of the cast alloy ingot. 6. A method as claimed in any preceding claim in which the alloy contains a total of 8 to 30 atomic percent of one or more rare earths. 7. A method as claimed ia any preceding claim in which the alloy contains 2 to 28 atomic percent of boron. 8. .A method as claimed ia any preceding claim in which the alloy contains ae least 50 ©comic percent of iron. 9. A method as claimed in any preceding claim in which the alloy contains praseodymium. 10. A method as claimed in any preceding claim in which the alloy contains one or mora of the elements AX. « Ho and Si. 11. A ©ethod as claimed ia any preceding claim in which the alloy contains 50 or less atomic percent of Co. 12. A mathod as claimed in any of claims 1-6 in which the composition of the alloy is substantially that of any of Alloys Nos 1-31 - 13o A method of making a permanent sua,gnat comprising melting and casting a Ce^P^ gFe^B^ alloy; hot working the cast alloy ingot at a hot working temperature of at o -4 least 500 Cs and at a strain rate de/dt ia the range of 10 to 10 per second where e is the logarithmic strain and fc is time; and forgoing the hot worked alloy ingot into a permanent magnet. 14. A method of making a permanent magnet comprising melting and casting an alloy whose constituents comprise at least one rare earths iron and boron5 hot working the cast alloy ingot at a hot working temperature of 1000°C and a strain rata -4 of de/dt in the range of 10 to 10 per second where £ is the logarithmic strain and t is time; and forming the hot worked alloy ingot into a permanent magnet. 15. a ma thod according to any one of claims lf 13 and 14 of making, a permanent n&grjefc'substdrrtlsLLly- as described in any of theTScamples. 16. A permanent magnet when made by the method claimed in any preceding claim. F. R. KEIIiY & CO. AGEMS FOR THE APPLICAMIS.