Chapter 2 Rare earth-cobalt permanent magnets

chapter 2 RARE EARTH-COBALT PERMANENT MAGNETS

K.J. STRNAT University of Dayton Dayton, Ohio USA

Ferromagnetic Materials, Vol. 4 Edited by E.P. Wohlfarth'~ and K.H.J. Buschow © Elsevier Science Publishers B.V., 1988 131

CONTENTS 1. I n t r o d u c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1. B a c k g r o u n d , d e f i n i t i o n s a n d s c o p e o f article . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2. Selecting c a n d i d a t e m a t e r i a l s f o r p e r m a n e n t m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3. E a r l y w o r k o n m a g n e t i c r a r e e a r t h - t r a n s i t i o n m e t a l alloys . . . . . . . . . . . . . . . . . . . . . . . 1.4 H i s t o r i c a l o u t l i n e of p r a c t i c a l m a g n e t d e v e l o p m e n t . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2. F u n d a m e n t a l s o f r a r e e a r t h - c o b a l t m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. P h y s i c a l m e t a l l u r g y a n d c r y s t a l s t r u c t u r e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Basic m a g n e t i s m o f r a r e e a r t h - t r a n s i t i o n m e t a l c o m p o u n d s . . . . . . . . . . . . . . . . . . . . . . 2.2.1. S p o n t a n e o u s m a g n e t i z a t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.2. C u r i e t e m p e r a t u r e . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.3. M a g n e t i c a n i s o t r o p y - P h e n o m e n o l o g y . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.4. M a g n e t i c a n i s o t r o p y - T h e o r e t i c a l c o n c e p t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.5. T e m p e r a t u r e v a r i a t i o n of m a g n e t i z a t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.6. S u m m a r y o f b a s i c m a g n e t i c p r o p e r t i e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. R e a l m a g n e t s f r o m R - C o alloys - G e n e r a l a s p e c t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1. I n t r o d u c t o r y r e m a r k s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.2. B e h a v i o r o f p o w d e r s a n d c o m p a c t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.3. C h e m i c a l s t a b i l i t y p r o b l e m s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.4. T y p e s o f m a g n e t i z a t i o n b e h a v i o r . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.5. M i c r o s t r u c t u r e s o f R - C o m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.6. P h a s e r e l a t i o n s in lhe S m - C o - C u s y s t e m . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.7. M a g n e t t y p e s b y a l l o y c o m p o s i t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3. M a g n e t m a n u f a c t u r i n g t e c h n o l o g y . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. M a g n e t f a b r i c a t i o n m e t h o d s - A n o v e r v i e w . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. C o m m e r c i a l a l l o y p l o d u c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. M a n u f a c t u r e o f s i n t e r e d m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.1. I n t r o d u c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.2. P o w d e r m i l l i n g a n d b l e n d i n g . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.3. M a g n e t i c a l i g n m e n t a n d c o m p a c t i o n . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.4. S i n t e r i n g a n d h e a t - t r e a t m e n t . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.5. M a c h i n i n g , h a n d l i n g a n d m a g n e t i z i n g . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4. B o n d e d m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.1. S l a t e of d e v e l o p m e n t . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.2. F a b r i c a t i o n m e t h o d s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.3. S u b t y p e s a n d t h e i r p r o p e r t i e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.4. M o l d i n g , m a c h i n i n g , h a n d l i n g a n d m a g n e t i z i n g . . . . . . . . . . . . . . . . . . . . . . . . . .

132

134 134 136 137 139 142 142 146 146 147 148 151 152 153 155 155 156 158 159 161 167 170 174 174 175 178 178 180 180 181 184 186 186 187 188 189

K.J. S T R N A T

4. P r o p e r t i e s o f c o m m e r c i a l m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. G e n e r a l c o m m e n t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. M a g n e t i c p r o p e r t i e s of s i n t e r e d m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. M a g n e t i c p r o p e r t i e s of b o n d e d m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4. M a g n e t i c s t a b i l i t y - T h e r m a l c y c l i n g a n d a g i n g . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.5. S u m m a r y o f p h y s i c a l p r o p e r t i e s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5. A p p l i c a t i o n s o f the r a r e - e a I t h p e r m a n e n t m a g n e t s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References ..............................................................

133

190 190 192 194 196 202 202 204

1. Introduction

1.1 Background, definitions and scope of article Since 1966 a new family of magnet materials has been evolving that is generically known as the "rare earth permanent magnets" (abbreviated as REPM in this chapter). Their engineering properties quickly surpassed those of all earlier magnets, with the best laboratory values for two common figures of merit, the room-temperature energy product and the intrinsic coercive force, each now being 5 - t 0 times greater than" the highest values for premium magnets of other types. * This is illustrated in fig. 1, which shows a chronology of magnet development milestones in this century. It should also be noted that both, (BH)max and MHo, have now exceeded 100 times the values of the best steel magnets available before 1930. Given such outstanding properties, the REPM have opened up entirely new device uses for permanent magnets (see, e.g., Ervens 1982a, Brandis et al. 1980, Strnat 1979), although they are - and will probably remain - too costly for universal use. Together with the hard-magnetic ferrites, however, they are rapidly broadening the field of applications for permanent magnets in general as well as displacing more conventional magnet types in many cases. As a commercial product, rare earth magnets had their start in the mid-1970s, about 20 years after the hexaferrite (ceramic oxide) magnets. The REPM are metal magnets. Their magnetically active components are alloys of one or more of the 3d-transition metals (abbreviated T M or T) with elements of the rare earth group (symbols RE or R). The rare earths comprise the fifteen 4f-elements, La (atomic number 57) through Lu (71), and Y (39). Of these, the elements included in fig. 2 are of particular significance for magnets in combination with cobalt as the main alloying component. We presently distinguish two major subgroups of REPM: the rare earth-cobalt magnets, most commonly based on S m - C o alloys, and the RE-Fe-based magnets typefied by N d - F e - B . The latter are a fairly new, "third generation" variety of the REPM, introduced in 1983, which * For definitions of the concepts and magnetic quantities used here (the primary system of units in this chapter is CGS) and discussions of different magnet materials see other articles in this H a n d b o o k (Zijlstra 1982, St~iblein 1989, Buschow 1988),~ also general textbooks on magnets (e.g., McCaig 1977, Burzo 1986). The textbook by Burzo has an exhaustive bibliography on the REPM, current as of 1985. 134

R-Co PERMANENT MAGNETS

135

kJ/m3 n'400

MGOe

50

(BH)mox k0e 6O

4O 300

Nd-Fe-B

30

50 MH¢

SIn-Co, Fe,TM "2-17"

R-Co "1 -5"

f

200

MA/rn

40

SmC°5~"F- 4. R-Co

"/

"1 -5"

[_] 3

"2-17" ] 30 Sm-Co,Fe,TM. . . . .

2O

2 f ALNICOS 1~ --

IO~

E

STEELS F RI 0 fgo01920 194019601980 YEAR

100

20

Nd-Fe-B--

I0

FERRITES~../--ALNICOS~ 7 I STEELS ~ . x I

I

;-:/

%0,920,940,960,980 0 YEAR

Fig. 1. Chronology of magnet development since 1900: Best reported laboratory values for the static energy product and intrinsic coercive force. (After Strnat 1986.)

have great commercial promise and are now in a period of intense study and rapid development. This chapter is only concerned with the RE-Co-based magnets (except where comparisons are in order); the R E - F e - B magnets and their derivatives are the topic of a separate chapter (Buschow 1988). The rare earth-cobalt magnets are discussed here with a fairly applied bias. Relevant fundamental subjects are treated exhaustively-elsewhere in this handbook series: basic magnetic properties of the RE-elements by Legvold (1980), those of R E - T M intermetallics by Buschow (1980), and the theory of hard-magnetic behavior - small-particle magnetization reversal, domain-wall nucleation and motion, and the role of anisotropy in these - by Zijlstra (1982). Buschow (1988) discusses the various mechanisms thought to be controlling the coercivity in all REPM, and the PRINCIPAL CONSTITUENTS

MINORITY CONSTITUENTS

Mo~. Lowe st/oSst'~ Moderate CostJ

HighestC:st "

Fig. 2. The rare-earth elements now used in rare earth-cobalt permanent magnets, their effect on magnet properties (referenced to Sm alloy), and relative cost of the rare earth component.

136

K.J. S T R N A T

theoretical models advanced in attempts to explain the multiplicity of observed magnetization reversal phenomena. We review here only general concepts regarding the origins of coercivity and its complex relationship to crystal anisotropy and to the microstructure of real magnets. The body of technical and scientific literature concerned with the R E P M has become so large that it is impossible to do it justice in the space available; many important publications must go unmentioned. In addition to "milestone papers", the selection made here favors review articles that integrate many contributions and where the reader can find references to the original articles. There exist a number of proceedings books of conferences entirely or to a significant extent devoted to the REPM and their uses, and some special issues of company journals. They are listed below to facilitate systematic background reading, and also to shorten the citations of individual papers contained in the collections. These collective works are attributed to their editors. An early monograph on Rare Earth Permanent Magnets was written by Nesbitt and Wernick (1973). A (continuing) series of International Workshops on Rare Earth Permanent Magnets and their Applications has so far produced eight books - of the second through the ninth Workshop (the first was not published). These workshops were edited by Strnat (1976); Strnat (1978a); Kaneko and Kurino (1979); Strnat (1981); Fidler (1982); Pan, Ho and Yu (1983); Strnat (1985); and by Herget and Poerschke (1987), respectively. A complementary series of Proceedings of Symposia on Coercivity and Anisotropy exists. The first Symposium was not published, the second was separately bound (Strnat 1978b), while the third and fourth such Symposia are included in the Workshop books of 1982 and 1985; the fifth was again a separate volume, edited by Herget et al. (1987). Mitchell (1985) edited the proceedings of a special European Communities Workshop on Nd-Fe Permanent Magnets. The company journals mentioned above are: Goldschmidt Informiert, No. 35 (Kornfeld 1975) and No. 48 (Kornfeld 1979); and Thyssen Edelstahl Technische Berichte 6 / 1 (Brandis et al. 1980).

1.2. Selecting candidate materials for permanent magnets In terms of basic magnetic properties, a ferro- or ferrimagnetic substance is a candidate for permanent magnet use when it fulfills the following three conditions (discussed in CGS system terminology): (1) It must have a high spontaneous magnetization, Ms, in the temperature range of practical interest (typically around room temperature). The saturation intrinsic induction, B i s = 4~rMs (in gauss) sets the upper limits for the remanent flux density, B r = B i s , and for the energy product, (BH)max = ( ~1B i s ) 2 (in MGOe). (2) The Curie temperature, Tc, of the main phase must be high enough for the contemplated application. The lowest Tc values now accepted by design engineers are about 300 ° C ( N d - F e - B and M n - A I - C magnets); the Tc = 460 ° C of hexaferrites qualifies them for many more applications; while the Curie points of R E - C o and Alnico type alloys, in the range from 500 to 900 ° C (but typically 700-800 o C), make theSe especially suitable for elevated-temperature use.

R-Co PERMANENT MAGNETS

137

(3) Finally, there must be a mechanism for creating a sufficiently high intrinsic coercive force, MHo. What is "sufficient" depends again on application requirements, but for modern magnet materials it is usually defined as MHc = ±B2is, the minimum value that permits a close approach to the theoretical energy product limit, (2~rMs) 2 = (~Bis) 1 2• Iron, cobalt, and especially F e - C o alloys have the best combinations known of Bis (at 20°C) and Tc, but no way has been found to give them a very high coercivity. F e - C o is the active component of Alnico and of the elongated-single-domain particle (ESD) magnets. Hc-values of 4 to 6 kOe were once expected for the latter; but the best coercivity that could in fact be derived from single-domain behavior, using shape anisotropy as the energy barrier, was about 2000 Oe (160 k A / m ) , or only about 30% of the value needed to realize the predicted energy products of about 50 MGOe ( - 4 0 0 M J / m 3) (see, e.g., Luborsky 1966). Incidentally, the highest conceivable room-temperature energy product is about 150 MGOe, for fully dense F e - C o with Bis = 24.5 kG, at least based on the present knowledge of ferromagnetism. A reasonable objective for further PM development aimed at the highest possible (BH)max is to strive for a metallic magnet material whose main phase is close to an F e - C o composition with 30-50% Co, in which secondary phases occupy only a very small volume fraction, and which has a minimum MHc commensurate with the Bis of the alloy, i.e., > 12-13 kOe. Magnetocrystalline anisotropy offers a better way than particle shape to generate sufficiently high coercive forces. Ferromagnetic crystals having a single strongly preferred axis of magnetization, and when in the form of micron to submicron-sized "single-domain particles", should require fields about equal to their anisotropy field, H A = 2 K / M s, for magnetization reversal. This was predicted by early fineparticle theory (Stoner and Wohlfarth 1948) and asserted by later refinements (see, e.g., Zijlstra 1982, or Livingston 1973). The 1950s and 1960s brought, therefore, a systematic search for candidate substances having a strong uniaxial crystal anisotropy. The hexagonal ferrites were the first group of such materials discovered to find commercial application; the RCo 5 group of intermetallic compounds was next, leading up to the broader family of rare earth magnets. 1.3. Early work on magnetic rare earth-transition metal alloy s

Until well after World War II the individual rare earths in metallic form were almost unavailable and of no industrial interest. From 1900 until the late 1950s there were only a few studies of rare earth-transition metal alloy systems reported. They used mostly Ce as the RE component and concentrated on the existence of intermetallic compounds and their crystal structures. Of the phases of interest here, CeCo 5 was discovered by Vogel and Frilling (1947), and Ce2Co17 by Zarechnyuk and Kripyakevich (1962). These were assumed to be prototypes for other, isostructural R E - C o compounds, which was later largely confirmed. The significance which the rare earths have in nuclear fission reactions then caused the development of a technology for the separation and reduction of the RE. This made the elemental metals commercially available and brought on a surge of scientific interest in RE

138

K.J. STRNAT

alloys and their properties. Many R E - T M compounds were prepared and systematic studies of their magnetic properties began. Wallace (1973), in a book, summarized and critically reviewed the large amount of work done in the preceding 15 years. Several investigations of the magnetic ordering phenomena found in R E - T M combinations yielded preliminary information that proved important for the later development of permanent magnets. They were done at Bell Telephone Laboratories where Nesbitt et al. (1962) studied G d - C o alloys; at the University of Pittsburgh where Wallace and co-workers prepared many RCo 5 and related Fe and Ni compounds (see, e.g., Nassau et al. 1960); and at the US Naval Research Laboratory. There, Hubbard et al. (1960) studied G d - M n and G d - C o alloys; they reported that GdCo 5 has uniaxial crystal anisotropy and that a powder of this compound had a coercive force of 8 kOe (0.64 M A / m ) . While this was indeed a first indication of the promise that the RCo 5 compounds held for permanent magnets, it was then generally ignored, probably because the low saturation of the specific compound, GdC%, disqualifies it as a practical magnet material. The work of all these groups was in part motivated by the hope to find new technologically interesting magnet materials that would ferromagnetically combine the large magnetic moments of the heavy RE and 3d-transition metals while retaining the high Curie points characteristic of Fe or Co. (One can, e.g., speculate that DyCo 5 might have a magnetic moment of 18.4/x B per formula unit, corresponding to a saturation induction of about 25.5 kG, if it were a simple ferromagnet in which all atomic moments add. If this were possible, we might now have permanent magnets with energy products approaching 160 MGOe.) However, these hopes were disappointed. The heavy RE consistently couple their moments antiparallel to the T M moments, so that DyCo 5 has indeed only a moment of 1.6/z B per formula unit, or about B s = 2200 G. Based on the above-mentioned early publications, several other laboratories joined the systematic study of the magnetism of R E - T M intermetallics. Among them were groups at CNRS Grenoble (James et al. 1962, Lemaire 1966); the USAF Materials Laboratory (Strnat et al. 1966a) together with the University of Dayton (Ray et al. 1964); the Technical University of Vienna (Kirchmayr 1966); and Philips Research Laboratories (Buschow and Velge 1967). While the interest was primarily in basic magnetic properties, it was necessary for all to study the binary R E - T M phase diagrams and the crystal structures of the many intermetallic phases which were unknown or poorly described up to that time. A review of early results (Weik and Strnat 1965) revealed a few among the binary R E - T M compounds that combine sufficiently high Curie point and saturation values with a low crystal symmetry, so that they might have potential as new permanent magnet materials. These were the RCo 5 phases in which R is a "light" RE (Ce, Pr, Nd, Sin) or Y, i.e., one of the RE elements having small or near-zero atomic moments, b u t which do couple ferromagnetically with the T M moments. Some R2Co17 phases were soon added to the list (Ostertag and Strnat 1966, Strnat et al. 1966b, Lemaire 1966), and so was LaCo 5 (Velge and Buschow 1968). Systematic studies of the magnetocrystalline anisotropy using single crystals (Hoffer

R-Co PERMANENT MAGNETS

139

and Strnat 1966, Tatsumoto et al. 1971, Ermolenko 1973, Ermolenko et al. 1973) as well as oriented powders (Strnat et al. 1967, Buschow and Velge 1969) showed that YCo 5 and most other RCo 5 compounds have an extremely large crystal anisotropy with a single easy axis of magnetization. This led to the prediction (Strnat and Hoffer 1966, Strnat et al. 1967) that these were indeed candidate materials for outstanding new permanent magnets. At the same time, the finding of an unfavorable easy-basal-plane anisotropy for Y2COlv (Hoffer and Strnat 1966) discouraged consideration of the 2-17 alloys for magnets until the R2Co~7, with R = Sin, Er and Tm, were also shown to be strongly anisotropic with an easy axis (Ray and Strnat 1972, Schaller et al. 1972, Narasimhan et al. 1974). Simple fine-particle hypotheses had suggested that micron-size powders of any RC% alloy, and of the latter three R2Co17 , should have very high coercivity; aligned compacts of such powders should be good magnets. However, the practical realization of useful REPM was not so simple and encountered many obstacles. When powders were produced by mechanical grinding, H c reached only 5-10% of the anisotropy field * (see, e.g., Strnat 1967). The coercivity so obtained was adequately high only for SmC%, and so the initial magnet development concentrated on this compound. Over a period of two years, fine-particle S m - C o magnets with increasing energy products were prepared in several laboratories, by means of epoxy bonding (Strnat et al. 1967) and by binder-free compaction (Buschow et al. 1968), but their magnet properties deteriorated slowly in air.

1.4. Historical outline of practical magnet development The development of a liquid-phase sintering technique by Das (1969) and by Benz and Martin (1970) made fully dense and stable SmCo 5 magnets possible. This method became the basis for the "first generation" of commercial REPM. At the same time, an American and a Japanese group discovered independently that the partial substitution of Cu for Co in SmC% (Nesbitt et al. 1968) and in CeC% (Tawara and Senno 1968) allowed the formation of a nonmagnetic precipitate in the RCo 5 matrix. This could impede domain-wall motion and cause usefully high coercive forces without the need for comminution. It was first thought that this discovery had opened a way for processing RC% into magnets by casting, similar to Alnico; but the procedure proved to be too complex. Copper-containing, precipitation hardened magnets did become commercial products, but they are now all processed by powder metallurgy, like SmCo 5. As the understanding of the alloy systems and the magnet processing technology improved, it became possible to use as the RE-component other elements, such as Ce, Pr, Nd and La - singly or in combinations that include the natural blend of these elements, known as mischmetal; but to consistently obtain high He, it was necessary to retain a substantial proportion of Sm in the mixture (see, e.g., Benz and Martin 1971, Nagel et al. 1976, Higuchi and Ishigaki 1979). In the Cu-containing magnets it was also possible to * In this chapter Hc is usually used to mean the intrinsic coerciveforce, MHc or Hoi. The induction coerciveforce is always abbreviated 8He.

140

K.J. STRNAT

replace some of the Co by Fe and to increase the total amount of transition metals beyond the stoichiometric 1-5 ratio to R(TM)5+x, where x < 1 (Tawara and Senno 1973). The realization of practical 2-17 magnets proved even more difficult and took another decade. All the magnet fabrication methods that had worked well for SmC%, when applied to Sm2Co17, yielded only coercive forces in the 1-3 kOe range, which is too low. Experimental work in the early 1970s demonstrated that the formation of quasi-binary intermetallics of the type R2(Co1_xFex)17 was possible for most rare earths. The introduction of some iron also stabilizes the easy-c-axis anisotropy in all the light-RE systems (except with Nd) and increases the saturation, while depressing the Curie temperature only slightly (Ray and Strnat 1972, Schaller et al. 1972, Yajima et al. 1972, Hamano and Yajima 1978). With this knowledge it was possible to predict a "second generation" of RE-Co-based magnets that could have energy products in excess of that for SmCo5, up to about 60 MGOe (Strnat 1972b). However, it remained difficult to obtain sufficient coercivity, even with R = Sm which here, too, yields higher anisotropy and greater He-values than the other rare earths. Over a period of several years, many different experimental approaches were tried, but they produced only marginally useful He-values, and energy products that were lower than those of RC% magnets. It was demonstrated that small additions of still other 3d-metals could be beneficial; yet serious practical problems persisted. The search of those years for useful R2Co17-based magnets was chronicled by Strnat and Ray (1975a), Tawara and Strnat (1976) and by Ervens (1979). The most pragmatic approach - very systematically developed by Tawara and co-workers was to extend the range of R(TM)z compositions that could be magnetically precipitation hardened with the help of Cu; this was possible to about z = 7.2 (Senno and Tawara 1974). The result was two-phase sintered magnets in which the main phase has the 2-17 structure. Careful heat-treatments made coercivities in the 4-10 kOe range possible for different compositions. Substituting Ce for some of the Sm, and Fe and/or Mn in the TM position, a whole family of commercial magnets was developed, all characterized by general wall pinning at homogeneous precipitates. Then Tawara and Senno (1976) were able to obtain high coercivity in sintered Sm(Co,TM)8-magnets containing some Fe and Cu; and Nagel (1976) achieved energy products in excess of those for SmCo5 (up to 30 MGOe at H~ = 12 kOe) with sintered magnets of the general composition Sm(Co, Fe, Mn, Cr)8.5. These first true 2-17 magnets all exhibited "nucleation controlled" magnetization reversal, i.e., their magnetization behavior is more like that of SmC% than that of R(Co, Fe, Cu)5+z. The Mn,Cr-containing magnet has the severe disadvantage that its coercive force drops very quickly with increasing temperature above 20 ° C. A very important step forward was the discovery that the addition of small quantities of zirconium, in Sm(Co, Fe, Cu, Zr)~, coupled with a relatively complex heat-treatment, yielded coercivities of 6.5 kOe (520 k A / m ) by general wall pinning in alloys with more Fe, less Cu and higher z-values than previously possible (Kaneko et al. 1977, Ojima et al. 1977, Yoneyama et al. 1978). The best energy

R-Co PERMANENT MAGNETS

141

product was also over 30 MGOe (240 k J / m 3) at a Zr content of 1.5 wt.% and z = 7.4. It is important to note that these magnets have a much smaller (negative) temperature coefficient of MHo than the Mn,Cr-containing alloys mentioned above, making them much more useful for elevated-temperature applications, where 2-17 magnets with their high Curie points were expected to excel. In was soon shown that small additions of titanium to similar alloys having z = 6 to 7, and with R = Ce or Sm (Inomata et al. 1977, 1978), and the presence of hafnium in Sm(Co, Fe, Cu, Hf)7.4 (Nezu et al. 1979) have similar beneficial effects on microstructure and coercivity. Then it was found that much higher coercive forces, 10 to > 25 kOe, were also possible, using slightly more Zr (up to 3 wt.%) and longer heat-treatments; the best energy product was raised to 33 M G O e (263 kJ/m3), combined with H c = 13 kOe, for an alloy having z = 7.67 (Mishra et al. 1981). Finally, Shimoda et al. (1979), also working with Sin(Co, Cu, Fe, Zr)z alloys, for use in polymer-bonded magnets, were able to raise H c up to 26 kOe at z = 8.35, and to 8 kOe even for an alloy with z = 8.94, which is on the TM-side of the 2-17 (z = 8.5) stoichiometry! What all the successful " 2 - 1 7 " magnets have in common is a very small-scale microstructure within the main grains. It is visible only in the electron microscope and consists of separated cells of the 2-17 phase, surrounded more or less by a thin shell of a 1-5 boundary phase, the crystal lattices being coherent. It is this boundary phase that pins the magnetic domain walls. Such a cell structure was first discovered by Livingston and Martin (1977) in a Sin(Co, Fe, Cu)~ magnet. Zr, H f or Ti additions promote the optimal formation of this cellular precipitate structure. Because of its great practical importance in generating high coercivity, this microstructure, the conditions of its formation, and its interaction with domain-wall motion have since been extensively studied by many scientists and on a variety of magnet alloys. For recent reviews, see the articles by Livingston (1986) and Ray (1986a). The REPM offer a possibility - unique among permanent magnet materials - to alter the temperature dependence of the remanence within wide limits by alloying measures. For some applications it is important to minimize the temperature coefficient of B r around room temperature. Such temperature compensation can be accomplished by partial substitution of, say, Sm in a 1-5 or 2-17 alloy with a heavy RE of the group Gd, Tb, Dy, Ho and Er. This was first practically demonstrated by Benz et al. (1974), and by Jones and Tokunaga (1976), who substituted Gd and Ho into SmC%. The potential of using Tb, Dy or Er was explored by Martis et al. (1978), and a great variety of such HRE-containing magnets were actually sintered by Narasimhan (1981). Li et al. (1980) prepared the first temperature-compensated 2-17 magnets of the low-Hc variety, Mildrum et al. (1983) applied the concept to high-H c 2-17 magnets by incorporating Er. The work on this problem was reviewed by Strnat and Tauber (1983). The R E - C o alloys for processing into magnets were initially always prepared by melting together the constituent metals. From the economic point of view it was very important when calciothermic reduction methods were later developed which combined several steps in the production of alloy powders and permitted the use of

142

K.J. STRNAT

rare earth oxides as starting materials, instead of the expensive metallic RE. Cech (1974) introduced a "reduction-diffusion" (R-D) method using Call 2 as the reductant. Domazer (1974) developed a similar, so-called co-reduction process which employs calcium metal vapor as the reductant and puts in a part of the TM-component also as an oxide. These methods significantly reduced the alloy prices at the time of their introduction. However, the rationalization of induction melting has kept pace and it continues to be extensively used. Herget (1975) published a thorough analysis of these alloy production methods. Finally, the development of bonded, or matrix versions of the REPM deserves attention. The first SmC% magnet samples prepared in laboratories were epoxy-matrix magnets, and commercial polymer-bonded magnets based on SmCo5 were introduced early (Taylor and Wainwright 1976); but they proved insufficiently stable for most uses. However, bonded permanent magnets in general have great technological and economic appeal, and bonded REPM clearly had great potential (see, e.g., Stmat et al. 1976a,b). So their development slowly continued, and by now it has resulted in a broad spectrum of commercial or near-commercial products employing a variety of magnet alloys, binders and production methods. The use of precipitation hardened 2-17 alloys has greatly enhanced the stability, and so has the use of better polymers or of soft metals as matrix materials. The technological importance of bonded REPM is now rapidly growing. New injection-molded, extruded and calendered products are being introduced, and the availability of N d - F e - B in addition to Sm-Co-based alloys has raised the prospect of much cheaper high-energy matrix magnets. (There are, however, new and more serious stability problems associated with the use of the N d - F e - B that need to be solved.) 2. Fundamentals of rare earth-cobalt magnets

2.1. Physical metallurgy and crystal structures The commonly used designations " R E - C o " magnets, "Sm-Co", and "1-5" and "2-17" for their main subcategories, are quite simplistic. The real magnets are multi-phase metallurgical systems with complex microstructures, generally not in an equilibrium state, and they always contain more than just two elements. Even the distinction between RE-Co and RE-Fe-based magnets is becoming increasingly blurred. However, it was the study of binary rare earth-transition metal alloy systems that provided the impetus for the REPM development. The metallic phases present in the magnets are all derivatives of binary intermetallic compounds between a 4f and a 3d-transition metal. For the understanding of the REPM it is therefore essential to first review the alloying behavior of selected binary R E - T M combinations, typical crystal structures of the compounds, and their relevant basic magnetic properties. Alloy of samarium with cobalt (or with Co and Fe) are technologically the most important. Figure 3 shows the Co-rich portion of the Sin-Co equilibrium phase diagram, basically after den Broeder and Buschow (1972) and Perry (1977), but with a modification of the liquidus line (shown dashed) according to Ray (1986b), who

R-Co PERMANENT MAGNETS

1500 [*c]

50

40

I

Wt. % Sm

I

I

.-" "

1300

/

20

.Y I

1400 ~

143 0

I

/1:310

°

1200 I100 1000

900 805 =

800

r-

~

8

700 30

20 At.%Sm

I0

0

Fig. 3. The Sm-Co equilibrium phase diagram. After den Broeder and Buschow (1972), Perry (1977) solid lines; Ray (1986b) dashed lines and 1310 ° C peritectic. found that Sm2COl7decomposes peritectically on heating and that the previously reported nearby eutectic does not exist in an equilibrium diagram. The R E - C o diagrams with other rare earths are generally similar, but with some minor, systematic variations. There is almost no terminal solubility in either the RE or Co component; but a number of intermediate phases exist (7 in Sin-Co). These are usually shown as line compounds, although certainly SmCo 5 and Sm2Co17, but also the other thoroughly investigated 1-5 phases (such as those of Ce and some heavy RE), have finite homogeneity ranges at higher temperatures. Figure 4 (Strnat 1972b) summarizes the binary R E - C o phases which have been reported to be stable at room temperature, or which can be easily retained there in a technologically useful metastable state by quenching from a higher temperature. For each idealized phase composition listed the crystal symmetry is given (in some cases for room- and high-temperature structures), and the bars indicate with which RE elements the phases form. For magnets, only those in the upper half of the bargraph are of interest: a 2-17 or 1-5 as the principal flux-producing phase; the others, down to 1-3, as secondary minor phases that influence the coercive force and its temperature dependence. These compounds are ferro- or ferrimagnetic above r o o m temperature. The two meandering dashed lines define the regions where the Curie

144

K.J. S T R N A T

:OMPOUND

,1

,._oIcel I,,,0Is,,,I GdI"b I'"'1 o, I'o l':r I "ml'u

RCol3

CUBIC

R 2 COl7(I)

HEX RHOMB

RC05 (2)

HEX

R 5 COl~ 3}

b0..I A

\\~r"-n R2Co 7 RCo 5

P

o od v

SYMMETRY

//Ji/f

.

,\\\\\\\\'I\\\\\\',,\\\\\\\\\\

Ti//J//f

el/IlIA . . . . . .

_

J

I

\"

HEX RHOMB RHOMB HEX RHOMB

~

CUBIC

RCo 2 L--J

R2Co 3

ORTHORHOMB

R 2 C°l .7

HEX

n," 2:)

R4 Co 3

HEX

OE

Rx=2Co (4)

I.iJ

I-. _w n,, 0

~A

V///////////A HEX

R24 COil R9Co 4

V-/A

R3Co (5)

V////////////////////M VAORTHORHOMB

ORTHORHOMB

Fig. 4. Summary of rare earth-cobalt intermetallic phases. (After Strnat 1972b). Notes: (1) Where double bars are drawn, the upper bar indicates the crystal symmetry and existence range of a high-temperature modification; the lower bar represents the structure reported to be stable at room temperature. (2) For the heavy rare earths, Tb and D y through Tm, the phase with the CaCus-type structure exists only as RCos+ x with a Co-excess which increases with increasing atomic n u m b e r of R. The RCo 5 phases are generally unstable at room temperature (see fig. 5), b u t can be retained metastably at room temperature by rapid cooling. (3) Additional 5-19 phases have been reported, but their existence - certainly as a stable room-temperature phase - is questionable.

point is above 20 ° and 400 ° C, respectively. Several peculiarities should be discussed: No stable 2-17 compound exists with La; instead, a cubic LaCo13 forms (Velge and Buschow 1968). The 1-5 phases of the HRE, Dy through Tm, exist only off-stoichiometry (RCo5+x) with a Co excess, x, that increases with increasing atomic number of R. Also, most 1-5 phases are stable only at elevated temperatures and decompose during cooling in a eutectoid reaction (Buschow 1972, 1974). The temperature ranges of phase stability are wide for the light RE but become quickly narrower with increasing atomic number of R. The solidus temperature, T~, and decomposition temperature, Ta, are shown in fig. 5. Small amounts of impurities on Co sites can significantly affect phase stability: - 3% Fe raises Td between 100 and 150 o C, narrowing the range of existence, while 3% A1 lowers Ta and stabilizes the 1-5 phase (Buschow 1974). With R = La, Ce, Pr and Nd, a stable 5-19 compound exists between 2 - 7 and 1-5 (Schweizer 1972, Ray et al. 1973, Khan and Feldmann 1973, Ray 1974). In

R-Co PERMANENT MAGNETS

145

T e m l )Q

['c I 1500

L melting

1000

~lon 500

0

X=A[ J r ~ , t ~ B : ~ ~ I r , L o C e Pr Nd Pm S m E u Gd T b Dy Ho Er Tm Yb Lu R - Element

Fig. 5. Stability ranges for the "RCo 5'' phases showing solidus temperature, Ts, and eutectoid decomposition temperature, Td. (After Buschow 1974). Note the pronounced effect on the phase stability of small third-element substitutions for Co (x = 3% Fe or A1).

addition, a metastable La5Co19 and a variety of secondary metastable phases, such as R C 0 4 , R 7 C 0 2 9 and R 4 C o 1 7 have been found to form easily with R = La, Ce, Pr and Nd in this composition range. These are part of a theoretically expected infinite homologous series of crystal lattices that converges toward RCo 5 (Strnat and Ray 1975b), The crystalline structures of the phases from 1-3 to 2-17 stoichiometry are closely related. All exist in a hexagonal form, and all but the 1-5 also have a rhombohedral modification. As the most important examples, fig. 6 illustrates the RCo 5 structure, of CaCus-type; t w o R2Co17 structures are shown in fig. 7: the hexagonal one of Th2Nilv-type (right), and the rhombohedral of ThzZn17 type

OCOBALTO RAREEARTH Fig. 6. Basic crystal structure of R E - T M compounds near the 1-5 stoichiometry (CaCu 5 type).

146

K.J. STRNAT

.~{j-

O COBALT

•

RARE EARTH

Fig. 7. Basic crystal structures of RE-TM compounds near the 2-17 stoichiometry rhombohedral ThzZn17 (left) and hexagonal ThzNil7 type (right). (left). Drawn are large, non-primitive cells that make the c o m m o n hexagonal symmetry elements obvious. Recently, still other modifications of the 2 - 1 7 structure have been identified in sintered magnets (Ray 1986a). Ray also states that the hexagonal form exists only with excess cobalt, as RCo x with up to x -~ 12. It is in crystals like these with a unique crystallographic axis that a high uniaxial magneto-crystalline anisotropy might be expected. And indeed, the c-axis in these structures is magnetically unique - in some cases the easiest and sometimes the hardest direction of magnetization. It was the discovery of a strong easy-axis magnetic anisotropy in G d C o 5, YCo 5, SmCo 5 and some related compounds that pointed the way to the R E P M (Hubbard et al. 1960, Hoffer and Strnat 1966, 1967, Strnat et al. 1967).

2.2. Basic magnetism of rare earth-transition metal compounds 2.2.1. Spontaneous magnetization Of the binary R E - T M compounds, only those rich in the 3d-metal Co or Fe (2-17 and 1-5) - and among these particularly the compounds of the low-moment, light lanthanides Ce, Pr, Nd, and Sm, and also of the nonmagnetic La and Y - have a sufficiently high spontaneous magnetization to be of interest for magnets. Their saturation induction values at 2 0 ° C are summarized in fig. 8. The previously discussed general rule for the 4 f - 3 d exchange interaction is evident: the light RE couple ferromagnetically with Co or Fe, yielding high saturation, while the heavy RE couple antiparallel, which results in a very low saturation for the compounds of

R-Co PERMANENT MAGNETS 16

i

i

I

I

i

i

i

i

i

i

i

i

147 i

i

1.6

14

1.4

M2I

1.2

I0

=£8

//

1.0

I

.J

I\1 ~ 1

/'1/I

6

0,8 t-

0,6

0,4

2

0

0,2

r

i

La Ce Pr Nd Pm SmEu Gd Tb Dy Ho Er Tm Yb Lu Y 57 58 59 60 61 62 63 64 65 66 67 68 69 70 71 59

0

Fig. 8. Room-temperature spontaneous magnetization values (" saturation intrinsic induction") of transition metal-rich RE-TM(-B) compounds of interest for permanent magnets.

the high-moment HRE, especially G d through H o (Nassau et al. 1960, Nesbitt et al. 1962, Lemaire 1966, Strnat et al. 1966a,b). Also included for comparison is a line (dashed) for the ternary compounds RzFea4B (after Sagawa et al. 1984), which obeys the same trend.

2.2.2. Curie temperature Magnet alloys must also remain ferromagnetic up to reasonably high temperatures, say 300 ° C or more. Figure 9 summarizes the Curie points of the same families of R E - T M compounds. Obviously, the R 2 C o ] 7 , with Tc = 800-950 ° C, are best. The RC% Curie temperatures are comfortably high (except for CeCos), while the R2Fe17 Curie points, all < 180 o C, are too low for practical magnets. However, the introduction of some boron to form the RzFe14B compound raised Tc b y 200-300 ° C over the corresponding RzFe17 values, which helped to make several of these phases - notably those with R = N d and Pr - potentially useful magnet materials. As a general rule, the compounds with cerium have much lower B s and Tc than their neighbors with La or Pr. (Since Ce tends to be tetravalent and therefore nonmagnetic, while transferring an additional electron to the 3d-shell of the TM-atom.) CeCo 5 with its Tc < 400 ° C, was long considered unsuitable for R E P M because its magnetization drops off too rapidly on heating. Now, however, the low T c = 3 0 0 ° C of NdzFe14B is accepted for m a n y appfications, suggesting that CeC%-based magnets may yet find a more prominent commercial position.

148

K.J. STRNAT 1200, oc

.

.

.

.

LOCOl3

°K

/-"~ 800

,..~ ~

,

k

R2CoI.-.f J,__~ ~._. ,_.._.

17~ 1200

I

RCo5 .~ ~

=

----4 ....., ~._

/ \ ~/

~

~

, ~ . . . ~ - - - - ' - ' ~ ~'--"

.~ ~

~

~

~

~

~

1000

~ . , L. ~ _ ~

600

R2FeI7

400

~

~

200

-200

-273

. . . . . . . . . . . . . . . . . . . . . . . . . . . ~ Lo Ce Pr Nd Pm Sm Eu Gd Tb Dy Ho Er Trn Yb Lu Y

0

Fig. 9. Curie temperatures of transition metal-rich RE-TM(-B) compounds of interest as the main phase in permanent magnets.

2.2.3. Magnetic anisotropy - Phenomenology

According to modern concepts of permanent magnetism, a high intrinsic coercive force can be expected of a ferromagnet that has a very large magnetocrystalline anisotropy, especially one with a single preference axis for the spontaneous magnetization vector, M S. High H c may then be achieved in several alternative ways, depending on the micromagnetic processes that dominate the magnetization reversal domain nucleation, localized or homogeneous wall pinning, or coherent spin reversal in single-domain particles (see Zijlstra 1981, Livingston 1973, 1986). There are theoretical models which predict that in either of these cases MHc could be as high as the anisotropy field, H A, which thus becomes an important quantity to know. In any case, the crystal anisotropy is the third basic magnetic property of primary importance for the R E - T M compounds. One way to describe it is by magnetization curves measured on crystals, with the H-field applied along different axes. Figure 10 shows pairs of such curves for four high-anisotropy compounds that have become very important as "parents" of REPMs: SmC%, 8 m 2 C o 1 7 , a 2-17 phase with Co + Fe, and Nd2Fe]4 B. In each case saturation is achieved in a very low H field applied along the c-axis, the "easy axis" (EA), which is also the direction preferred by M s in the absence of a field. By contrast, it takes a very high field to saturate a crystal in any direction in the basal plane, perpendicular to c. We speak of a " h a r d plane", which the magnetization vector avoids. (There can, however, be differences between directions within that plane, a "basal-plane anisotropy".) Note that only two of the curve pairs, those for SmC% and Nd2Fea4B, were indeed measured on individual single crystals; the other two are for powders consisting of many small -

R-Co PERMANENT MAGNETS

L

61 IIc-AXIS

1 [kG] tl

/

"

149

ND2FE14B

?i-:

i.s.2c017

.....

/ 6 / Z O

Z

0

,00

2O0

E X T E R N A L FIELD, H

[kOe] 3OO '

Fig. 10. Easy- and hard-axis magnetization curves of several high-anisotropy compounds on which practical permanent magnets are based.

particles, most of them also crystals, whose easy axes were oriented parallel by a magnetic field and fixed in that position. Measurements on such samples can only yield information about the main, axial components of the anisotropy, but not on any variations in the basal plane. However, since larger single crystals of new substances are often difficult to produce and were initially unavailable, and since it is mostly the uniaxial anisotropy component that determines how useful a compound is for magnets, m a n y such measurements on powders have been made (see, e.g., Strnat et al. 1967, Buschow and Velge 1969, Ray and Strnat 1972.) The field strength at which saturation is achieved in a hard direction is called the anisotropy field, H A. It is a theoretical upper limit for the coercivity. An alternative description for the anisotropy uses constants, K; in a commonly used simplification, often just the first-order anisotropy constant, K 1. For the uniaxial magnetic symmetry described above, the (cgs) relationship then holds: H A = 2 K J M s, where M S is the spontaneous magnetization. Ka can be interpreted as an energy density, namely, the work per unit volume needed to rotate M s through 90 o from the easy c-axis into the basal plane. K 1 c a n also be negative; the c-axis is then the hardest axis to magnetize and the magnetization vector prefers directions in the base plane. This generally makes magnetization reversal easy; the coercivity of such "easy basal plane" (EBP) materials remains low, and thus they are not suitable for permanent magnets. Finally, there are also intermediate preferred directions possible, which lie on an "easy cone" (EC), a situation that can allow high H c as long as the cone angle remains small. The mathematical description of this requires additional, higher-order anisotropy constants. As the temperature changes, transitions between the described anisotropy types can occur ("spin reorientation"). For good PM materials one wants to have EA anisotropy with a high H A in the intended use temperature range (usually around normal room temperature). Practi-

150

K.J. STRNAT X

R2C°17

0 L

~ 1

1

~ 1

1

~ 1

1

1

[

~8 1 1

; M

~////////////////////~

%c57 5 ]

Fe

Ce2C°t7

V///////~

I

t

Nd2COl7 V / / / ~ / / / / / / / / / / / / / / / / / / A

%c0,7

~/////////////A

I M~

%%,

V//////////~

I A~

s.:~% 7 s%%7 t---1

I al

V///,////////////A Gd2C°~7 V//~ o,2%, V / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / / A

F,

Er2C°I7 I

Fe

E~Co,7

I

~'/////////////////A

IMn

I

E~ Cot7 t~-f//////////A

I

rr": % 7 f

Fe

I Ni

~///////////////A

Fe

Y5% T~C°17

5c~,7 5co17

Fe

[~ V/////////////////A V//////////////////////~

] Ni

V//A~I

] AI

Y2Calf

] C.

Y2C°17 P ' / / / / / A

[7/-A easy plane

~

easy axis

m/xed phase OF unknown

Fig. 11. Influence of transition metal substitutions on the anisotropy (magnetic symmetry) of R2Colvbased intermetallics of compositions R2(C%_~Mx)l> (After Hamano and Yajima 1978, Ray and Strnat 1972.) cally all of the binary RCo~ phases show this desirable behavior. Only TbCo51 has an EBP, while N d C % has a very weak EA anisotropy at room temperature and develops an EC on cooling below about 20 o C. A m o n g the binary R2C%7 phases EA behavior is rather the exception, shown only by the compounds with R = Sm, Er and Tm. However, it was found that relatively small substitutions of certain third elements, especially iron, for some of the cobalt can induce the EA anisotropy in most other R2COlv , except those where R = N d or Dy. Figure 11 shows qualitatively how a partial or full replacement of Co by Fe, but also partial substitution by Mn, Ni or A1, affects the anisotropy of such 2 - 1 7 phases. It should be noted that isostructural binary 2-17 compounds exist with Co, Ni and Fe, but not with Mn or A1. The latter elements therefore tend to destabilize the 2-17 structures, and only small amounts can be substituted for Co.

R-Co PERMANENT MAGNETS

151

2.2.4. Magnetic anisotropy - Theoretical concepts

The nature and the physical origins of the high crystal anisotropy in R E - T M alloys have been extensively investigated during the last decade. Alloys which can serve as the principal phase in permanent magnets, namely the hexagonal 1-5 and 2-17 compounds with T = Co or Co + Fe, and more recently also the R E - F e - B phases of 2 - 1 4 - 1 stoichiometry, have received special attention. Systematic experiments in which composition and temperature were varied have revealed a rich variety of spin-order and moment-orientation phenomena. For comprehensive reviews of the RCo 5 and R2(Co, Fe)l v behavior see the articles by Asti and Deriu (1982), Ermolenko (1982) and Wallace et al. (1982). Scientifically motivated experimental studies involving the partial substitution of third and fourth elements for RE and TM laid the basis for understanding the factors which affect anisotropy, and thus for manipulating them to achieve desired effects. It is now fairly well known which substitutions promote EA anisotropy or increase K 1 and HA, or how one can shift the temperature ranges where spin reorientation occurs. The theory of anisotropy has by now also evolved to a point where all the various contributing mechanisms have been identified, and where reasonably satisfactory quantitative calculations are possible for some of the factors involved. However, there are so many contributing effects, and their mathematical description is so complex and difficult, that there is little hope of "designing new or better magnet materials" using these theories. In any case, up to now, all technologically useful innovation has resulted from experimental work, guided by enlightened intuition, with theoretical explanation following behind. The magnetic structure of the R E - T M compounds is usually described as consisting of two sublattices comprising the RE and the T M atoms, respectively, each being ferromagnetic within itself. (Atoms of the same species in non-equivalent crystal lattice positions may, however, have different magnetic moments.) The net moments of these two sublattices add for the light-RE components, while they are antiparallel and thus subtract for the heavy RE. Regarding the anisotropy one can consider, in a first approximation, that the RE and T M species interact separately with the crystal lattice, the total magnetic anisotropy of the compound being the sum of the individual sublattice anisotropies. At higher temperatures, approaching the Curie points of the RCo 5 or R2(Co, Fe)17, the RE moments are largely disordered and the strongly ordered TM Sublattice determines the overall anisotropy behavior. (This is the reason why all RCo 5 have EA at high temperatures.) At low temperatures approaching 0 K, the anisotropy of the RE sublattice makes itself strongly felt (except, of course, for the nonmagnetic rare earths La, Ce 4+, Y and Lu). Different RE favor either c-axis alignment (in the RCo 5 case, Sm and Er), or basal-plane orientation of their moments, as do Pr, Nd, Tb, Dy, and Ho in RCo 5 (Greedan and Rao 1973, Buschow and van Diepen 1974.) In the first case, the Co anisotropy is strengthened (SmCos), in the second it is reduced and in some cases overwhelmed at cryogenic temperatures (NdC%, TbCo 5), leading to a spin reorientation toward the basal plane on cooling, with an intervening easy-cone range. The anisotropy of the RE sublattice is considered to be due to the influence on the localized R 3+ ion of the electronic field of its local atomic environment ("crystal

152

K.J. STRNAT

field"), and to that of the molecular field ( R - C o exchange interaction). A single-ion anisotropy model is satisfactory for the mathematical description. The nature of the T M anisotropy, particularly also that of the Co sublattice in RC%, is less clear, with the 3d-electrons having been variously treated as localized or as itinerant by theoreticians. However, it seems that in the RCo 5 the Co-sublattice anisotropy can also be fairly well described by a single-ion model with a large orbital component of the Co-ion moment being responsible for the strong EA anisotropy. In 2-17 compounds, the atoms located in the "dumb-bell sites" make a strong contribution to the anisotropy. These sites are preferentially occupied by the substituents when Co is partially replaced by other T M atoms, and this can even change the sign of K 1, inducing EA behavior (Deportes et al. 1976). This effect has also been described in terms of a single-ion, local crystal-field model; but certain results seem to require that band-structure changes also be considered (Perkins and Str~issler 1977) i.e., that 3d-electrons be treated as collectivized. The net anisotropy of the compounds is determined by the individual sublattice anisotropies and the inter-sublattice exchange interaction. In RC%, these anisotropy and R - C o exchange energies are of comparable magnitude, which has the consequence that high external magnetic fields can significantly distort the magnetic structure from the simple model of collinear moments (Ermolenko 1982). This is supposedly not so in the R2Co17 , where the R - C o exchange dominates over the crystal-field interaction (Wallace et al. 1982). Some observed anomalies in the temperature dependence of K 1 and the hard-axis magnetization curve of R E - T M crystals have been explained by a moment-canting model. This takes into account deviations from the strictly collinear magnetic order previously assumed, due to competition between the sublattice anisotropies (Ermolenko 1976, Rinaldi and Pareti 1979, Groessinger and Liedl 1981).

2.2.5. Temperature variation of magnetization In either of the families of " R E P M parent compounds", 1-5 and 2-17 (and in some others as well), two basic forms of the temperature dependence of spontaneous magnetization are found (see, e.g., Wallace 1973, p. 146ff; Buschow 1980, p. 338ff). They are shown qualitatively in fig. 12, using SmCo 5 and GdC% as examples. The ferromagnetic compounds of Sm and of the light rare earths (LRE) show typical Brillouin behavior, i.e., M s declines monotonically as the temperature increases, toward zero at Tc. (Some earlier reports of "dips" in this curve at low temperatures can be attributed to lack of saturation in an insufficient measuring field, due to the extremely high anisotropy. Spin-reorientation phenomena in some compounds can also complicate the behavior.) The compounds of the high-moment heavy rare earths (HRE), however, are ferrimagnets. When the magnetization of the H R E sublattice at absolute zero is greater than that of the Co-sublattice (or, more generally, of the T M sublattice), M s versus T curves have the general shape of the curve labeled GdCo 5 in fig. 12, with a compensation point. For the R E sublattice with its weak, indirect R - R exchange interaction, the magnetization drops off much more rapidly with T than for the strongly coupled Co sublattice. At T~omp the moments cancel, and at higher temperatures the Co (or TM) sublattice moment

R-Co PERMANENT MAGNETS

153

USE RANGE ,Ir~Ms, LRE-Co (M's) sMco5)

O.7M~" , X . . ~ LREI_xHREx-CO(MS) iI ', \ . . \ "

'..I J

\'.\..~

MS"~

_

0 TCOMP TROOM TEMPERATURE [K]

' i HRE-Co(MS) (E.o, GDCo5

TCURIE

Fig. ]2. Temperature variation of the spontaneous magnetization in LRE-TM and H R E - T M alloys (qualitative), and principle of the internal temperature compensation in REPM. (After Strnat and Tauber 1983.)

dominates. Above the compensation point there is a temperature range in which M s rises with increasing T. This has the important practical consequence that an internal temperature compensation is possible in R E - C o magnets by mixing (at least) one L R E and one H R E in the alloy. One can play out against each other the positive and negative slopes of the two curves in a limited temperature interval by the judicious choice of the H R E and its amount in the alloy, x, and thus make M s fairly independent of T by placing the M s maximum in the middle of the anticipated use-temperature range of the magnet. If the coercivity of a sintered magnet is high enough, its remanence will fairly closely follow M s versus T. This compensation procedure is qualitatively illustrated in fig. 12 for the example of Sm0.vGd0.3Co5.

2.2.6. Summary of basic magnetic properties Table 1 is an attempt to compile numerical values of the relevant fundamental magnetic properties for the phases which form the basis of R - C o permanent magnets. It includes the binary RCo 5 with Y and the light RE, also Sm2Co17 , and various multi-component modifications of these that are representative of commercial magnet types. The approximate theoretical upper limit of the energy product is also given, calculated from the room-temperature saturation magnetization. It has in fact been fairly closely approached (75-80%) by the best laboratory magnets of SmCos, PrC% and (Sm, Gd)C%, while there is no hope to come close to the (BH)ma~ potential of NdCo 5, given the low room-temperature anisotropy, and therefore He-limit, of this compound. It should be noted that many of the figures given are still somewhat uncertain for the reasons indicated in the footnotes to the table. The numbers for mischmetal (MM) alloys are particularly questionable since the exact composition of this RE blend varies from supplier to supplier, and different researchers have also used the term MM to indicate somewhat different

154

K.J. STRNAT

TABLE 1 Basic magnetic properties of compounds of interest for RE-Co magnets [room-temperature values (except Tc) ]. Data for single-crystal and powder samples. The experimental data are taken mostly from the following review papers: Ervens (1979, 1982a,b), Narasimhan (1986), Nagel (1980), Kirchmayr and Poldy (1979), Strnat (1972b); and references cited therein. Compound

Bs (kG)

YCo 5 LaCo s CeCo5 [3] PrCos NdC% SmCo5 [4] MMCo5 [5] MM0.sSm0.2Co5 [5] Sm0.6Gd0.4Co5 Sm2C%7

K 1 [1]

Tc (T)

(K)

( o C)

10.6 1.06 903 630 9.1 0.91 840 567 7.7 0.77 653 380 12.0 1.20 893 620 12.2 1.22 910 637 11.4 1.14 1000 727 - 9 . 0 - 0 . 9 0 - 7 6 8 ~495 - 9.8 - 0.98 - 773 - 500 7.3 0.73 1000 727 12.5 1.25 1193 920 Sm 2 (Coo.7Feo.3) 17 14.5 1.45 1113 840 Sm2 (Coo.8Feo.lMnoA)l 7 13.1 1.31 am(Coo.s7 Cu 0.13) 7.8 10.9 1.09 1120 847

(107erg/ cm3 [6]) 5.5 6.3 6.4 8.1 0.24 11-20 -6.4 -7.8 7.7 3.2 3.0 4.3 3.3

H A [11 (kOe)

Theor. (NO) max [21

( M A / (MGOe) m)

130 10.3 175 13.9 210 16.7 170 13.5 5 0.4 250-440 20-35 -180 -14.3 -200 -16 264 21.0 65 5.2 52 4.1 81 6.5 77 6.1

28.1 20.7 14.8 36.0 37.2 32.5 -20.2 ~24.0 13.3 39.0 52.6 42.9 29.7

(kJ/ m3) 223 164 117 286 295 259 ~160 -190 106 310 417 340 236

Notes: [1] In most cases H A was measured as the hard-axis saturation field and K 1 calculated from H A = 2K1/M

s.

[2] The limiting energy product values were calculated as (½Bs)2 [CGS]. [3] A wide range of Bs and Tc values has been reported, possibly due to the use of impure cerium. Since the most likely RE-impurities (La, Pr and Nd) increase Bs and To the lowest reported values are listed. [4] For SmCo5 a wide range of property values is found in the literature, especially for the anisotropy (HA, K1). Some discrepancies are attributable to the fact that most experimenters had only fields much lower than H A available and made a long extrapolation. The use of powders instead of single crystals also brings inaccuracies. However, "SmCo5'' exists off-stoichiometry (as SmCo5+x) when quenched in from higJaer temperatures, and H A varies significantly with x. And there is indeed some confusion about the correct definitions of K 1 and H A. [5] The Ce-rich RE mixture "mischmetal" (MM) is produced with a fairly wide compositional range. Thus, properties reported for MMCo 5 vary widely depending on source and composition of the MM, especially the Ce content. Unfortunately, the MM composition was often not reported. [6] 107 erg/cm3 = 106 J / m 3.

c o m p o s i t i o n s . T a b l e 2 shows the c o m p o s i t i o n ranges f o u n d in typical c o m m e r c i a l "Ce-rich mischmetals" from various sources (Wells and Narasimhan

1979). T h e

c o m p o s i t i o n of a specific lot d e p e n d s on w h e t h e r b a s t n a e s i t e or m o n a z i t e minerals were used, the source of these ores, the c o n d i t i o n s u n d e r w h i c h the c h l o r i d a t i o n a n d metal-reduction steps were conducted, and whether elements might have been added t o d i s c o u r a g e c o r r o s i o n . T h e u s e o f M M f o r m a g n e t s calls f o r a c l o s e c o o p e r a t i o n b e t w e e n t h e R E s u p p l i e r s a n d m a g n e t m a n u f a c t u r e r s , a n d it r e q u i r e s c o s t l y q u a l i t y control measures which may largely offset the perceived economic advantages of using this basically i n e x p e n s i v e a n d a b u n d a n t R E mixture.

R-Co PERMANENT MAGNETS

155

TABLE 2 Composition ranges for typical Ce-rich mischmetal. RE component

wt%

Lanthanum Cerium Praseodymium Neodymium Other rare earths

20-30 46-55 4-7 13-20 1-3

Typical impurities Iron Oxygen Ca, Mg, A1, Si, C may be present dependingon container and electrode materialsused in processing

wt% 0.1-4 0.1-0.3

Note: Several percent of magnesium is sometimes added to commercial mischmetal to reduce the rate of atmospheric oxidation.

2.3. Real magnets from R - C o alloys - General aspects

2.3.1. Introductory remarks None of the intermetallic compounds identified above as potential hard-magnetic materials is per se a permanent magnet. Real magnets are more complex in several respects: the overall chemical composition of the magnet (or its main magnetic component) is not that of a stoichiometric compound; there are always several phases present; the metallurgical microstructure often exhibits intricate details (some of which are essential for getting high coercivity, while others are not) and it is typically not in thermodynamic equilibrium; and some magnets are composites that contain nonmagnetic binders as secondary components which magnetically separate discrete particles of the alloy. Also, the processes by which the magnetization state changes in the R E P M are much more complicated than those which fine-particle theory predicted. Considering this last aspect first: It is impractical, perhaps even impossible, to prepare single-crystal particles of a 1-5 or 2-17 compound that show single-domain behavior in the classical sense, reversing from full positive to full negative saturation in a single event at a very high field strength near HA; much less is it possible to make magnets of a useful size consisting of m a n y such inherently single-domain particles. The magnetization reversal in real particles, or in the grains of sintered magnets, never seems to proceed by a coherent spin-rotation process, but always by the movement of domain walls, or by the local formation of small reversed domains followed by wall motion. The variety and complexity of microstructures possible in the R E P M can provide several mechanisms for domain nucleation and for the pinning of domain walls. A good understanding of the factors which control the important structural features and their interaction with the wall motion is, therefore, essential for controlling the engineering properties of the magnets, especially the coercivity. Finally, the R E P M are never just binary alloys (with the possible exception of " S m C o 5''). Certain additional elements are purposely added to achieve special basic

156

K.J. STRNAT

magnetic properties (HRE for temperature compensation, Fe to raise the saturation, etc.); or for economic reasons (e.g., Ce, La, Fe); or to promote the formation of a desirable microstructure (Cu, Zr, etc.), Some nonmetallic impurities, primarily oxygen, are inevitable; but indeed, oxides may play a crucial role as domain-wall pins in some REPM (Ishigaki and Higuchi 1978, Schweizer et al. 1971). We shall now discuss these complexities and their consequences for the magnet properties in more detail.

2.3.2. Behavior of powders and compacts In all early attempts to make REPM under laboratory conditions (see, e.g., Strnat 1967), the near-stoichiometric R - C o intermetallics were prepared as coarse-grained ingots, which were then mortar ground or ball-milled into powders, yielding a high percentage of single-crystal particles. The particles were placed in a mold, a magnetic field was applied to saturate and orient them by aligning their easy axes, and they were then fixed in this position either with a resin binder or by compaction to a high packing density. While this simple approach provided the "proof of existence" of REPM, it did not result in practically useful magnets for the reasons mentioned below. The coercivity of various 1-5 and 2-17 alloy powders prepared by such mechanical grinding initially increases with decreasing particles size (or longer milling time), but then it reaches a maximum ( - 10-20 kOe for SmCos, but only between 1 and 4 kOe for YCos, PrCos, MMC% and Sm=Co17 ) and declines again on further grinding. Figure 13 illustrates this, using as an example the same MMC% alloy comminuted in two different mills. It also becomes more and more difficult to align the particles in a field. This is attributed to a progressive disruption of the crystal lattice by plastic deformation during the grinding which locally lowers the an5000

~000

COERCIVE FORCE, MHc, Oe 3000

ff~

r\

/ \

2000

1000

i I (Ce'MM)C°5 I I ,--250# 10

\

(Ce-MM) Co s ~-

I 20

30

40 50 0 2 4 MILLING TIME, hours

6

8

10

Fig. 13. Effect of milling time and method on the coercivity of aligned (Ce-MM)Co~ powders prepared in two different types of ball-mill: (a) Rotating-jar ball-mill; (b) Small laboratory vibration mill. (After Strnat 1967.)

R-Co PERMANENT MAGNETS

157

MHc Oe] I0,000

9,000

8,000

SmCo&sCuL3sFeo.4-precipitation heat treated []

O.,=

(4 hrs. at 475°C after homogenization)

7,000

6,000

SmC% o,-

Sm Co&5 Cul.3s Feo.4-homogenized at IlO0=C,quenched 5,000

I 0

i I0

I 20

I 30

I 40 AGING TIME

I 50 [HOURS]

I 60

I 70

Fig. 14. Deterioration of the intrinsic coercive force of several R E - T M alloy powders during short-term exposure to air at elevated temperature. (After Strnat 1970.)

isotropy, especially near the particle surface (Zijlstra 1970). However, the severity of these adverse effects depends on the milling method employed. It is possible to increase MHc several-fold subsequently to milling by removing the disturbed outer layer of the particles, either with acids or by reacting it with Zn and thus deactivating it (Becker 1969, Strnat and Tsui 1970). Stress relief by annealing after grinding can also increase MHc, and grinding in liquid nitrogen can yield higher coercivity, presumably because there is less plastic deformation and more brittle fracture at these low temperature (Strnat and Tsui 1970). Laboratory magnet samples with good initial properties were made from such powders milled to particle sizes in the 5-50 Fm range and compacted, with or without a binder, by uniaxial die pressing, isostatic compaction, or by a combination of the two (Westendorp and Buschow 1969, Strnat et al. 1967). However, the properties of these magnets, and of the powders from which they are made, were found to be unstable over prolonged periods of time, especially at higher temperatures in air (Strnat 1970, Suzuki et al. 1979, Evans et al. 1982, Cremer et al. 1982, R.M.W. Strnat et al. 1982.) The primary magnetic effect of such aging is a loss of coercivity that is initially quite rapid and severe, then continues at a lower rate (see fig. 14). It can be slowed, but not prevented, by compaction to high density or by

158

K.J. S T R N A T

various binders. Eventually, after longer exposure, the B,H-curve shape in the second quadrant and the remanence are also affected. All these observations are consistent with the basic assumption, implicit in several different models of the magnetization reversal, that small particles of the single-phase 1-5 and 2-17 alloys reverse their magnetization by the shifting of domain walls, initiated by "nucleation" events on their surface if the particles are first saturated. Sufficiently high magnetizing fields can either remove walls completely or compress residual reversed domains to tiny volumes in a disturbed lattice region where their walls are strongly pinned. It is in the surface layer of cold-worked particles where such wall fragments are most likely pinned. Conversely, easy renucleation of domains - often cited as the mechanism initiating the magnetization reversal - would also most likely occur on the surface in places where the anisotropy is lower, where large local demagnetizing fields may exist due to the surface topography, and where RE (Sm) depletion by oxidation may form soft-magnetic second phases such as Co. So it is not surprising that the "nucleation" fields, and thus Ho, depend sensitively on the surface condition of the powder particles. Semi-empirical models for these processes were developed, initially by Becket (1968) and Zijlstra (1971), on the basis of measurements on individual particles. However, moderately fine powders of the copper-modified alloys, when properly heat treated, do not suffer the described aging loss of Hc; nor do compacts made from them (Cremer et al. 1982, R.M.W. Strnat and Luo 1982). The reason is that their coercivity is caused by a homogeneous intra-granular precipitate that impedes the movement of domain walls wherever they are located. So, Ho does not strongly depend on the surface condition of these particles. If, however, the precipitate is dissolved by a homogenization heat-treatment, these alloys act just like SmC%. This is also illustrated in fig. 14, showing early experimental results on a S m - C o - C u - F e alloy near the 1-5 composition. It is equally true of the newer " 2 - 1 7 " alloys, which usually contain other minor elements, such as Zr, in addition to Cu.

2.3.3. Chemical stability problems Surface oxidation of particles was mentioned as one cause of the coercive-force loss during air aging of powders. However, oxidation (or other forms of corrosion) can cause stability problems in a more direct way as well. All rare earths have a high chemical affinity for oxygen. As a consequence, in the presence of air and particularly at high humidity a n d / o r elevated temperature, the RE near surfaces will progressively corrode and the alloy composition will change. This is, of course, especially severe for fine powders with their large exposed surface area; but compacts and even sintered bodies will also slowly oxidize. In extreme cases these can disintegrate after exposure periods of several months to years. The oxidation rate of massive bodies is aggravated when open pores or microcracks are present that connect to outer surfaces. The affinity for oxygen of the different RE elements varies significantly, with the light RE being less stable than the heavy RE and Y. Among the former, Sm, Nd and Ce act relatively much better than La and Pr (whose presence as an alloying constituent has been known to cause sintered magnets to age unacceptably fast and, eventually, to crumble). The relative oxida-

R-Co PERMANENT MAGNETS

159

tion tendency of several important magnet alloys has been characterized by measuring the Weight gain of powders and pellets with time (R.M.W. Strnat and Luo 1982). This ease of oxidation has several implications for magnet fabrication and use: It has necessitated the development of protective coatings for particles used in polymer-matrix magnets, used sometimes even on massive sintered magnets. There is also a special need for caution in large-scale powder processing. Particles in the micron size range are often pyrophoric; powder batches can be lost to spontaneous fires and there is a definite danger of accidental explosions in REPM production. Even when this danger has been minimized by proper precautions, it is necessary to protect fine powders from oxidation by proper handling and storage, or by fast processing, to avoid compositional shifts that will affect the magnet properties and the densification behavior during sintering. For sintered magnets, a near-theoretical density (i.e., low porosity), the avoidance of microcracks, and the minimizing of excess RE-rich phases in the grain boundaries are essential for good corrosion resistance and magnetic stability. This is even more important for alloys with a high content of La, Pr or mischmetal than it is for SmC%.

2.3.4. Types of magnetization behavior The magnetization reversal of the small crystals comprising real magnets - whether they are well-separated particles or grains in dense, sintered or cast bodies generally proceeds by domain-wall motion. It is customary to distinguish two basic behavior patterns and, accordingly, classify magnets as "nucleation controlled" (here called type A) or "pinning controlled" (type B). These can be identified by the characteristic shapes of their initial magnetization curves and by the dependence of remanence and intrinsic coercive force on the peak magnetizing field (fig. 15). However, it should be noted that these types represent limiting cases; the mecha-

TYPEA B-H

TYPE B

"

",.m/

H #

I ,~ /

IOk/

B-H

/ 0

IOk"

I~~j .A'FUR__AATION ___ VALUE

IH,. 50kOe

0

I

I H

I0 k

50 kOe

,.

MHc

~ l 0

i

I

I = Hm

I 50 kOe

..........~ H m

Fig. 15. Schematic magnetization curves for "nucleation controlled" magnets, type A, hke SmCos; and for " p i n n i n g controlled" magnets, type B, like Sm(Co, Cu, TM)x-

160

K.J. STRNAT

nisms are sometimes found mixed (type C), transitions between them can occur as the temperature changes, and the same magnet can show quite different behavior after different heat-treatments. It is nevertheless important to discuss the two extreme cases. First the basic concepts: Type A. The movement of existing walls within a given grain is easy, while the "nucleation" of a reversed domain after removal of the walls is difficult. Type B. Small reversed domains exist at all times (or they form easily in low demagnetizing fields), but defects present in most of the magnet volume strongly impede the further movement of walls. Type-A behavior is characteristic of binary SmCo 5 (in sintered or bonded-powder form), its derivatives containing other (additional) RE, and of certain copper-free 2-17 magnets (Menth and Nagel 1978). Figure 15A shows the virgin curve rising steeply; after thermal demagnetization there are many walls present in the interior of the grain where they can be moved easily through the undisturbed crystal lattice by a small driving field. When the field is reduced again from a low peak value, Hm, and then increased in the negative direction, the shape of the demagnetization curve traced ("minor loop") still indicates mostly reversible wall motion with low remanence and coercivity. As the value of H~ is increased, B r rises slowly and MHc very slowly; it requires very high magnetizing fields to fully develop the best possible "major" demagnetization curve. This is because an increasing number of grains are cleared of walls, and reversed nuclei must now be formed at the higher nucleation field in each such grain before its magnetization can be reversed again. (Most grains act independently of their neighbors.) The physical details of this so-called nucleation process are still unclear in spite of extensive experimental and theoretical efforts [see the articles by Fidler (1981), Kronmtiller (1982), and Livingston (1986) for comprehensive reviews]. One can argue whether (or when) a truly new nucleus is formed in a previously saturated grain, or whether the reversal starts from a highly compressed residual domain left from the prior opposite magnetization state. The term "nucleation" must certainly be broadened from the classical definition to include the latter case of a n c h o r i n g / unpinning wall fragments at highly localized sites in the grain boundary regions of sintered magnets (Schweizer et al. 1971, Strnat 1972a, Liu et al. 1982). Optical and electron-micrographs of the metallurgical microstructure in sintered SmC% magnets typically show additional small grains of a n S m 2 C o 7 phase, oxide particles and voids; but the best such magnets have mostly single-phase 1-5 grains and rather clean, thin grain boundaries (Fidler 1981). Type-B behavior is exhibited by " b u l k hardened" Cu-containing magnets whose compositions can cover the range from 1-5 to 2-17. The main R E component in these is typically Sm, Ce, or Sm plus Ce (Nesbitt et al. 1968, Tawara and Senno 1968, Chino et al. 1979, Inomata et al. 1979), and there can also be some Fe or Mn substituted for Co. Of particular technological significance are the low-coercivity versions of the " 2 - 1 7 " magnets containing Fe, Cu and a small amount of Zr, H f or Ti (Yoneyama et al. 1978, Nezu et al. 1979, Inomata et al. 1977). The materials show the described characteristics in the as-cast or sintered state, and also as coarse particles, provided they have been properly heat-treated.

R - C o P E R M A N E N T MAGNETS

161

%] OF SATURATION 8i, s 80

'

A

60

B

C

40 20 0

I

0

25

50

75

I00

Fig. 16. Comparison of the virgin magnetization curves for typical magnets of three basic behavior types: A is nucleation controlled; B is pinning controlled, and C is mixed (or pinning with very nonuniform pin strength). (After Ervens 1982a.)

Typical qualitative magnetization curves are shown in fig. 15B. The virgin curve is almost horizontal at first; then the magnetization builds up to near saturation in a narrow range of H m. When the field is reduced and reversed, there is again almost no magnetization change until the same field range, near MHc, is reached on the negative side. The coercivity and remanence of minor loops at first remain near zero as Hm increases, then they rise steeply to their respective "saturation" levels. Only an H~ slightly greater than MHc is needed to fully develop the second-quadrant curve. These magnets are thus easy to "charge". Responsible for this behavior is the presence of a fairly homogeneous precipitate inside the main-phase grains which impedes the motion of domain walls wherever they are located. For the precipitates to be effective pins they must by crystallographically coherent with the matrix, have basic magnetic properties (saturation, anisotropy or exchange energy) that differ significantly from those of the matrix phase, and they must have the right size and spacing. For comprehensive discussions of these relationship see, e.g., the articles by Livingston (1986), Durst and Kronmiiller (1985), Kronmiiller (1982), Fidler and Skalicky (1982), Menth and Nagel (1978), Barbara (1978), and Oestereicher (1984). As was mentioned before, there is now a newer precipitation-hardened magnet variety that shows still another form of the virgin magnetization curve. In fig. 16, which compares all these magnetization behaviors, this is shown as type C. It is observed in Sm(Co, Fe, Cu, Zr)x magnets with x = 7.2-8 (Mishra et al. 1981) and in similar alloys that have been heat-treated to a state of very high coercivity (12-30 kOe). The behavior of these magnets has been interpreted as a mixture of localized and homogeneous wall pinning (Durst and Kronmiiller 1985). Direct observations of domain patterns in variable fields (Li and Strnat 1984) suggest that the pinning strength of the intragranular precipitates varies strongly from grain to grain in these magnets, while it is quite uniform in the low-Hc version. 2.3.5. Microstructures of R-Co magnets The metallurgical microstructure clearly controls the coercivity of all R E - T M magnets. As a consequence, much effort has been invested in developing and using

162

K.J. STRNAT

techniques for studying the details of metallurgical and magnetic-domain structures; and a major thrust of any project to create or improve new magnets must be directed at favorably affecting the microstructure by proper alloying measures and heat-treatments. Unfortunately, magnet preparation and analysis are frequently not done in the same laboratories; the experience base of the individual investigators is often too specialized, and there is not enough communication, Thus, many of the observations reported and conclusions drawn must be taken with some caution. Extensive bodies of microstructural results with magnetic interpretations have been generated by several groups, the following being representative publications: Fidler et al. (1983), Hadjipanayis (1986), Rabenberg et al. (1982a), and Livingston (1986). The structural features which are effective in domain-wall pinning or nucleation are generally of such small dimensions that they can only be seen in high-resolution transmission electron microscopes (TEM). Identifying the compositions and crystal structures of the significant precipitates or grain boundary phases often requires (and sometimes still defies) the capabilities of today's best microprobe analysers. And lattice imperfections such as stacking faults, which may be effective as domain-wall pins or nuclei, can only be made visible by lattice imaging techniques. Such observations in electron microscopes cannot be made on the massive real magnets, nor can magnetized samples be studied, and the possibilities of simultaneously observing the important structural features and magnetic domain walls are quite limited, especially under conditions of varying applied magnetic fields and temperatures. Thus, our knowledge of the interaction between the metallographic features and the moving domain walls is still in large measure based on clever conjecture and on generalizations of theoretical models that are typically too simple. There are many differences of opinion and even contradictions found in published "explanations": e.g., precipitates of the same secondary phases (2-7 or 2-17 in a SmC% matrix) are variously considered to be effective domain-wall pins that cause high coercivity, or easy nucleation sites that reduce it. (It is possible for either phase to do both, but only under special circumstances.) Nevertheless, the characteristic microstructure types associated with the different important magnet categories and magnetization behavior patterns are now fairly well established and shall be described below.

SmCosype magnets The commercially most important REPM is still sintered "SmCo5''. This and the other magnets exhibiting type A, or "nucleation-controlled" magnetization behavior can have nearly single-phase, featureless microstructures. (Although there is in fact often an excess of Sm over the 1-5 stoichiometry present in the form of magnetic Sm2Co 7 grains and Sm203 particles.) Microstructural aspects important for the permanent magnet properties are the matrix grain size, a crystal texture, the nature of the grain boundaries, and the kind and distribution of any secondary phases. Grains must not be too large to keep low the probability that they contain a nucleating defect. (Typical grain size in SmCo5 is _<10 ~m.) An easy-axis texture is usually desired and is brought about by compaction in a magnetic field. (Although there are some isotropic REPM in production.) This parallel alignment of the c-axes

R-Co PERMANENT MAGNETS

163

of the grains can be made nearly perfect, so that the orientation of adjacent grains often differs only in the basal plane. This minimizes the disorder at grain boundaries and, together with the near absence of voids, oxides and other secondary phases in well-prepared magnets, it makes the boundaries thin and clean looking in TEM images. However, H c can be very high, so these boundaries must still block the movement of domains walls from grain to grain, or else the first nucleation event would cause the reversal of the whole magnet. Thus, we must assume that there are imperfections associated with such boundaries that are not visible in the microscope, yet quite effective in disrupting the exchange coupling between grains, or otherwise capable of strongly pinning domain walls there. Considering the small wall thickness ( - 5 nm in SmCo 5), extremely thin layers of secondary phases in the grain boundaries of these so-called single-phase magnets could well have such an effect. Intragranular precipitates are absent in SmCo 5 in its high-coercivity state. However, since the SmCo 5 phase itself is unstable below 750-800 ° C (den Broeder and Buschow 1972, Martin and Smeggil 1974), prolonged heating below this temperature brings about the dissociation of the metastable phase present in sintered magnets. Its products are small 2 - 7 and 2-17 crystallites, with at least the latter serving as easy nucleation sites. The magnetic effect is a drastic reduction of H c on annealing around 7 0 0 - 7 5 0 ° C (Westendorp 1970). An alloy on the Co-rich side of the homogeneity region, SmC% +x, initially maintained as a single phase by rapid cooling, will also precipitate out 2-17 upon heating (Khan and Qureshi 1973) with the same adverse consequence for H c. To achieve optimum coercivity, SmCo 5 magnets must be annealed after sintering somewhat above the eutectoid temperature, i.e., around 900 ° C, and then rapidly cooled to below 400 ° C or so to avoid dissociation. Annealing at higher temperatures yields lower H c again, which has been explained as a weakening of the pinning ability of the grain boundaries (den Broeder and Zijlstra 1976).

R(Co, Cu)5-type magnets Copper substitution for some of the cobalt in SmCo 5 and in CeCo 5 and proper heat-treatment was found to lead to moderately high coercive force for these alloys in bulk form (Nesbitt et al. 1968, Tawara and Senno 1968). The magnetic hardness is caused by a homogeneous precipitate inside the grains of the 1-5 phase which can impede domain-wall motion. The general magnetization behavior is type B - " pinning controlled". The Cu has several other desirable effects: it allows the co-substitution of Fe (which increases the saturation), it lowers the eutectoid decomposition temperature and thus stabilizes the 1-5 phase (which improves the aging stability of the magnets), and it broadens the homogeneity range of the phase toward higher transition metal contents (to about 1-7 in some cases). On the negative side the non-magnetic Cu reduces B r and (BH)max. The precipitate in a C e - C o - C u - F e sample was identified by Leamy and Green (1973) as very small (10 nm), coherent particles of 2-17 phase. Hofer (1970) reported for S m - C o - C u alloys a separation into a SmCo 5 phase with some Cu in solution and a Co-modified SmCu 5 phase. This was said to come about by spinodal decomposition around the composition of SmCo3Cu 2 where the coercive force

164

K.J. STRNAT

Fig. 17. Example of the precipitation microstructures found in 1-5 type RE-Co magnets bulk-hardened with copper: SmCo3.sCul.0Fe0. 5 homogenized at 1100 ° C/3 h and quenched. Left: SEM image after aging at 575 o C / 4 h. Right: Light micrograph after additional aging at 525 ° C / 4 h. The precipitates are platelets of "SmCu 5"(Co, Fe).

became extremely high. Both papers speak of needle-like particles. Kamino et al. (1973), Katayama and Shibata (1973), Perry (1977), Glardon and Kurz (1979) and others have since then studied the phase diagram of the quasi-binary SmCos-SmCu 5 system, some as part of broader investigations of the ternary S m - C o - C u phase ternary diagram. All confirmed that there is complete mutual solid solubility at high temperatures and a wide miscibility gap at room temperature, enabling the precipitation; but there is considerable disagreement in detail. Mildrum et al. (1970) found in SmCo3.sCUl.0Fe0. 5 samples that were overaged far beyond the point of maximum H c, a Cu-rich precipitate that had the form of thin platelets parallel to the base plane of the 1-5 crystal structure, see fig. 17. In all cases the coercivity is due to precipitated discrete particles with magnetic properties different from the matrix that can pin domain walls.

R2(Co, Fe, Cu, M)lz-type magnets Magnets in which the principal phase is 2-17 were also made technologically useful by a bulk hardening process that leads to type-B or type-C (pinning-dominated) magnetization behavior. All " 2 - 1 7 " magnets in commercial production now contain some iron to increase Bs, copper to permit the magnetic precipitation hardening, and often a small amount of another element, M = Zr, Hf, Ti or a mixture of these, to aid in the formation of the microstructure needed for the latter. The RE component is usually Sm, but in some products this is combined with Ce, Pr, Nd or

R - C o PERMANENT MAGNETS

165

Fig. 18. Example of the precipitation microstructure found in bulk-hardened 2-17 magnets of the high-He type: Sm(Coo.67Fe0.23Cu0.08Zro.o2)8.35 fully heat treated to MHc = 23 kOe. Left: section containing the c-axis; note the rhomboid cells of the twinned 2-17 matrix, the 1-5 boundary phase surrounding the cells, and the thin bands of "z-phase" in the base plane, crossing many cells and walls. Right: basal plane section. (Courtesy A.E. Ray.)

Y (because Sm is scarce) or with a heavy RE (for temperature compensation); it is also possible to make useful " 2 - 1 7 " magnets with Ce alone. These alloys cover a range of RE-to-TM contents, if the formula is written as RTx, of approximately 7 < x < 9. The following discussions pertains to " 2 - 1 7 " magnets of the most common variety, with R = Sm and having compositions in the range 25-27 wt.% Sm, 14-20% Fe, 5-10%, Cu 1.5-3% Zr, balance cobalt and oxygen. (The microstructures of the other 2-17 magnets are similar.) Their typical microstructure differs from that of the Sm(Co, Fe, Cu)5 magnets discussed before. Its main feature is a more or less well-developed network of very small cells of a matrix phase, within the much larger grains, which are separated and often completely surrounded by a thin boundary phase of 1-5 stoichiometry and CaCu 5 structure (Ray et al. 1987). Figure 18 shows as an example two transmission electron micrographs of a cast sample of composition 8m(Co0.67 Feo.23Cu 0.08Zr0.o2)8.35 heat-treated to a very high coercivity, MHo = 23 kOe. The section (a) on the left side contains the c-axis while (b) is a basal plane section. (a) shows the cells to have a roughly rhombic shape, elongated in the c-direction, while they have irregular cross sections in the base plane. When such high-Ho magnets are heat-treated to their optimal magnetic properties, the cells have linear dimensions of about 100-200 nm and the cell walls are typically 5-20 nm thick. The cell interior has the rhombohedral modification of the 2-17 structure (2-17R) and i's heavily twinned, with the

166

K.J. STRNAT

ISOTHERMAL AGING DISORDERED 2:17 R

DISORDERED"2:17" H

DISORDERED 2:17 R

Sm - poor

Sm - rich

Zr - rich

Zr - poor

(Platelet Phase)

1:5

ORDERED2:17 R

Fe - poor

Fe - rich

Cu - rich

Cu - poor

Zr - free

Zr - poor

(Boundary Phase)

(Cell Phase)

SLOW OR STEP COOLING Co,Fe StoiC0, Cu)54---1:5 ~ Cu

2:17 R - -

{ 2:17

H--"~"

"="

Sm2(Co, Fe)17 Sin(Co, Fe, Z r - V a c ) l O _ l Z

Fig. 19. Proposedisothermal aging reactions for disordered "2-17R" and subsequent diffusion directions during step aging. (After Ray 1986a).

twin boundaries in the base plane. There are also other, very thin layers visible that are parallel to the basal plane and run across many cells and cell boundaries. They are a third phase, the so-called "platelet phase" or "z-phase" which contains most of the Zr (or Hf, Ti in other magnets). These platelets are now believed to have the hexagonal 2 - 1 7 H structure. ( " 2 - 1 7 H " seems to exist only off-stoichiometry, transition metal-rich of 2-17.) All three phases are crystallographically fully coherent in good magnets. Worldwide, many research groups have studied the metallurgy and microstructure of bulk-hardened 2-17 alloys. The interpretation of the images presented here follows a recent thorough model analysis by Ray (1986a) of the stages through which this microstructure evolves during the aging heat-treatment of a previously homogenized alloy. It attempted to take into account the earlier work by others. Ray's concept is summarized in fig. 19 (see also fig. 32). All the features visible in micrographs are developed during the isothermal aging at a temperature around 800 o C. Yet, a sample quenched from this state has low coercive force! The high coercivity develops only during a subsequent slow (continuous or stepwise) cooling to about 400 * C, and it can still be somewhat improved by prolonged holding at 400 o C. However, no further changes can be seen in the TEM micrographs. Thus, whatever structural or compositional changes occur during the slow cooling that create the effective domain-wall pinning sites, they are subtle and on an extremely small scale. Since they cannot (yet) be experimentally identified, the details of the magnetic hardening mechanisms are wide open to speculation at this time. Metal-

R-Co PERMANENTMAGNETS

167

lurgists now seem to agree in general terms on the topography and crystallographic identity of the phases present in the magnets; but there is little unanimity as yet concerning the reactions by which these phases develop from the homogeneous high-temperature parent phase (a disordered 2-17R), or on the role which minor chemical constituents play in controlling these processes, and there is even less understanding of how the magnetic domain walls interact with the structural features. It should be noted that the microstructures of the low-coercivity versions of 2-17 magnets (with less Zr) are very similar to that shown in fig. 18; but all dimensions are smaller. Typical cell sizes are about 50 nm and the boundary phase is only 5-10 nm thick. Using the Fresnel or Foucault techniques with a transmission electron microscope it is possible to observe stationary magnetic domain walls in thermally demagnetized samples, superimposed on the microstructure (see, e.g., Fidler and Skalicky 1982, Rothwarf et al. 1982). Such images show that the domain walls run roughly parallel to the c-axis, as one would expect, but they follow the zigzag path of the 1-5 cell boundary phase. This has been interpreted in two ways: either, the domain-wall energy is higher in the 1-5, so that the boundary phase would repel a wall (Livingston and Martin 1977); or the domain wall has a lower energy in the 1-5 than in the 2-17 cell interior. This implies an attraction of the wall by the cell boundary (Nagel 1979), which seems a more logical reason for the walls to conform to the boundary shape in the absence of a magnetic field. In any case, it is the 1-5 cell walls which impede domain-wall motion; and because of the very small cell size, the effect on the magnetization reversal is more like homogeneous wall pinning (type B) than localized grain-boundary pinning (type A). In fact, it appears from Kerr-effect domain observations that in both high-Ho land low-He 2-17 magnets the grain boundaries are magnetically soft and provide easy nucleation sites (Livingston 1975). However, the domain growth in an applied field was shown to proceed quite differently in the two subtypes of 2-17 magnets (Li and Strnat 1984), and the high-Hc version exhibits strong inhomogeneities from grain to grain, magnetically as well as chemically (Rabenberg et al. 1982b). Therein seems to lie the reason for the different magnetization behavior: the low-Ho magnets are uniform and the precipitate provides a low and equal pinning strength everywhere; in the high-Ho magnets there is a wide distribution of grain compositions and therefore differences both in the precipitate structure and in its interaction with domain walls from grain to grain. 2.3.6. Phase relations in the S m - C o - C u system

The growing technological importance of the bulk-hardened REPM stimulated several systematic studies of phase equilibria in multi-component alloy systems containing copper. While it seems hopeless to try and establish the actual phase relations in the 5- to 6-component systems that comprise today's "2-17" magnet alloys, ternary phase diagrams are a step closer from the binary Sm-Co diagram to describing these complex compositions. The equilibrium diagram Sm-Co-Cu was of particular interest because it contains some of the early commercial REPM compositions. It could also serve as a model for other commercial alloys containing

168

K.J. STRNAT

single phase

Cu

twophase

cu

A

"

2o/~__~o /

f

sm

6o~_

20

\

/ \

At 1200°C

~

~o Atorn-°/o Co ~

so

SmcSumCU6

~

~

2o.,~ SmC0,/I'~

"~

~k'°

,

- - A t 800°C - -

/ eo I \ co Sm2C°? SmC°5 Sm2C°17

'SmCo ° / d3 C Y lSmCo5Sm2Co17°°

Fig. 20. Isothermal sections of the Sm-Co-Cu phase diagram. [From Herget (1979), after Perry (1977); modified by author.] additional rare earths as admixtures to Sm. Of equal importance for the newer bulk-hardened magnet alloys containing iron is the system S m - C o - F e , which was recently studied by Schneider et al. (1985). The Co-Cu-rich side of the S m - C o - C u system was first investigated by Perry (1977), and the work was later extended by Glardon and Kurz (1979) with a special view toward preparing magnets by directional solidification. Figure 20 [after Perry (1977), slightly modified by this author] shows two isothermal ternary sections: one at 1200°C (left) where only the cobalt-rich compositions are solid, and one at 800 ° C (right), just below the eutectoid decomposition temperature of the binary SmCo 5. Figure 21 shows three isopleths, quasi-binary temperature versus composition plots containing the 1-5 and 2-17 phase fields, with the Sm content as the compositional variable. The Cu content is kept constant at 0, 10 and 20 at.%, respectively. With this information it is possible to discuss the events leading to bulk hardening of the RE(Co, Cu)5+x alloys in some detail. Above Td ~ 805 ° C, the value of the SmCo5 eutectoid decomposition temperature reported by Perry, and up to the solidus, there is complete solid solubility between SmCo 5 and SmCu 5. At 800 ° C and below, "SmCos" with no or very little Cu sluggishly precipitates out 2-17 and becomes Sm-rich, with n o lattice coherency; this has the previously mentioned disastrous consequences for H c. However, the incorporation of Cu stabilizes the 1-5 phase, lowering Td and increasing the Co solubility at high temperatures. It also changes the lattice parameters of the 1-5 and 2-17 phases toward each other in

R-Co PERMANENT MAGNETS

169

I/,00

1200 o

b

I---

I/,00 r

I000

1200

P

00C

I--I~00 600

20

-115

1200

o

800

o

l--

600

20

15

10 800

6oo

20

~I--

15

~0

5

Co

Atom-% Sm

Fig. 21. Three isopleths from the ternary S m - C o - C u phase diagram at 0, 10 and 20 at.% Cu, respectively. (After Perry 1977,)

such a way that a coherent precipitation of 2-17 becomes possible (Melton and Perkins 1976), the mechanism of magnetic hardening described earlier by Leamy and Green (1973). On the other hand, at higher copper contents the 2-17 phase becomes unstable, and a low-temperature miscibility gap develops along the SmCos-SmCu 5 quasi-binary fine. It becomes possible to decompose a Sm(Co, Cu)5 phase into a Cu-stabifized "SmCos" matrix phase and a Cu-rich 1-5 precipitate (Hofer 1970, Kamino et al. 1973), creating the microstructures illustrated in fig. 17. Sm2Cu17 does not exist. As a consequence, copper is much less soluble in the Sm2Colv phase than it is in 1-5, and its solubility limit is strongly temperature dependent. (About 15% at 1200°C, 5% at 800°C, see fig. 20.) When alloys with Cu contents in this range are quenched and then annealed around 800 ° C, homogeneous precipitation of the coherent 1-5 network takes place which can cause coercivities of several kOe by wall pinning. Overaging causes the additional formation of a metallic Co (Cu) phase that is magnetically soft and destroys the coercivity again by easy domain nucleation. Of course, in practical magnet alloys iron is also present, plus some Zr (Hf, Ti), but the metallurgical mechanisms seem to remain basically the same. According to Ray (1986a), an important effect of Zr is that it stabilizes the 2-17R phase and

170

K.J. STRNAT

inhibits the precipitation of the soft Co(Fe, Cu)-phase, thus permitting the retention of 2 - 4 times higher Fe contents. It also appears to enhance the compositional differences between the 1-5 boundary phase and the matrix, thus increasing the wall energy difference between them and the pinning efficacy (see, e.g., Rabenberg et al. 1982a,b). The compositional modifications of real magnets away from the ternary S m - C o - C u alloys discussed above mean that details of the sohitionizing and aging treatments must be somewhat different. They should indeed be carefully optimized for each new magnet alloy. Specific examples of heat-treatment cycles that result in good hard-magnetic properties will be given in section 3. A more recent systematic study by Schneider et al. (1985) of the phase relations in the S m - C o - F e system will be an important aid in understanding and further developing magnet alloys of this kind.

2.3. 7. Magnet types by alloy composition So far in this chapter we have introduced many different RE-Co-based alloys with potentially useful magnetic properties, and we have discussed various metallurgical, structural and micromagnetic aspects of their hard-magnetic behavior. We shall now attempt a summary, reviewing the alloy types presently in production and their qualities, with some emphasis on technological importance and economic factors. The rare earth-cobalt based magnets now commercially available may be classified by compositional types as follows. (We shall begin with the "nucleation-controlled" alloys.) (1) " S m C % " magnets have a composition near, but on the Sm-side of the stoichiometric 1-5 ratio. The rare earth metal is either pure Sm or a less expensive rare earth mixture containing only - 7 0 % Sm. This " R E C o s " yields magnetic properties substantially identical to those of SmCo 5 (Domazer and Strnat 1976). Extremely high coercive force and knee-field values are easily achieved with SmC%. (2) (Sm, Pr)C% is a variation of the above with praseodymium substituted for a part of the Sm. Such magnets can have still higher energy products, although nobody has succeeded in achieving the superior remanence and energy product values that were predicted for pure PrCo 5. Praseodymium is somewhat more abundant than samarium and can help ease the chronic Sm shortage. Pr reduces the coercivity and can adversely affect the long-term stability. However, recent results obtained with oxide dispersions yielded SmCos-like coercive force and suggest improved elevated-temperature stability for sintered magnets in which 80% of the RE component is Pr (Ghandehari and Fidler 1985). (3) MMCo 5 magnets are magnets made from a base alloy in which the rare-earth component is either the conventional Ce-rich mischmetal (see table 2) or a modification thereof with reduced cerium content. The magnetic properties, Br, MHc, and (BH)max , are lower than for SmCo¢ and because of the relatively low Tc = 500 ° C and the greater oxygen affinity of Ce, La and Pr, the temperature stability is poorer; but it is still sufficient for many applications (see Mildrum and Wong 1976, Bachmann 1977). To keep the coercive force at a useful level some Sm is almost always added, typically 15-25% of the total RE content. Figure 22 shows demagne-

R-Co PERMANENT MAGNETS [kOe]

- 25

[kO]

15

I

I

- 10 I

-5 I

FI

MM1 x Smx Co5

9

765-

20

x = 0.4

432-

171

T/// //xoo7io

1-

0.9 0.8 -0.7 0.6 M 0.5 l -0.4 -0.3 -0.2 -0.l

I

[MA/m] -2.0

1.6

- 2.2

-0.8

-0.4

~H Fig. 22. D e m a g n e t i z a t i o n curves of sintered M M 1 _ x S m ~ C o 5 m a g n e t s with d i f f e r e n t s a m a r i u m contents.

(After Nagel and Menth 1975). tization curves of sintered M M 1 xSmxC% magnets with different levels of Sm addition, x = 0 to 0.4 (after Nagel and Menth 1975). In a more recent publication, Walkiewicz et al. (1983) reported that small alloying additions of magnesium together with copper can increase He to SmC% levels without using any Sm. Ce-rich mischmetal corresponds to the natural mixture of the rare earths in their most abundant ores, so it is indeed the cheapest and most abundant RE product available. When the use of SmCo 5 began to grow rapidly in the mid-1970s, threatening a samarium supply shortage which has since then become a serious impediment to broad commercial application of R E - C o magnets, M M - C o was seen as the principal R E P M product for the future. Extensive efforts were made to develop various MM-based magnet types to production maturity (see, e.g., Martin et al. 1974, Nagel et al. 1976, Wells and Narasimhan 1979), and a number of companies began manufacturing such magnets. However, a war-induced supply crisis in 1978/79 made cobalt a rather scarce and expensive commodity; so, to take best advantage of its magnetic flux potential, it was judged more economical to combine Co only with Sm. The cobalt price and supply have since recovered, but now there exists N d - F e - B which uses neither Sm nor Co and will be competing with M M - C o for the large-quantity consumer applications. These circumstances have so far kept M M C o 5 from finding its anticipated important place in the magnet market. (4) (Sm, HRE)Cos: This is the generic formula of a small family of magnets in which different amounts of G d or of other heavy rare earths are used to reduce the temperature coefficient of B r, even to near zero. Such magnets are needed in small volume for microwave tubes, measuring instruments, accelerometers, gyros and some other special devices where cost is secondary to critical performance requirements. They are very expensive compared with standard " S m C % " magnets. (5) Among the pinning-type magnets, Sm(Co, Fe, Cu)5_ 7 can have energy products comparable to sintered SmCo 5 while replacing some of the Co by cheap

172

K.J. STRNAT

Fe. Such magnets are also easier to magnetize. Ce(Co, Fe, Cu)5 has properties poorer than (MM, Sm)Cos; but it is made completely without Sm, using instead the most abundant of the individual rare earths, Ce, and it also contains some iron. Both of these magnet types have found significant market niches, particularly in Japan. With magnet alloys of this type it is also possible to combine Sm and Ce in any ratio, and this has allowed the development of a whole family of (Sm, Ce) (Co, Fe, Cu)5_7 magnets tailored to cost-performance compromises dictated by different applications (see Tawara and Chino 1979). (6) The precipitation hardened " 2 - 1 7 " magnets, particularly Sm(Co, Fe, Cu, Zr)x, with x = 7.2-8.5, are now rapidly gaining technological importance. They are produced in low- and high-Hc versions and are used in sintered and polymer-bonded form. The low-Ho magnets of this type are an extension of the Sm-based magnets discussed in (5), with similar applications. The high-H~ magnets have the potential of becoming a cheaper - and generally better - replacement for SmC% in all high-temperature applications requiring the best possible stability. However, they are so far more difficult to produce and to magnetize. (7) (Sm, HRE)TMx: These " 2 - 1 7 " compositions similar to those discussed in (6) above can also be temperature-compensated by replacing some of the Sm with one or more HRE elements. Magnets with H R E = Gd or Er are now in commercial pilot production. They will be used for the same special applications as the " 1 - 5 " magnets discussed in (4) and may eventually replace them.

STATIC ENERGY PRODUCT OF R E - TM MAGNET TYPES 0

50

I00

I

R Co5

150

I

I

200

250

I

I

V////////////M

R 2 (Co, TM )1? V//////////'//~////~

t

(BH)max

.300

350

I

I

Lk,J/m3.1 r

3

400

BONDED MAGNETS

- -_]

(Sm, HRE) Co 5 (MM, Sin)Co 5

v//////////A--2---_

SmCo 5

i

(Sm,Pr) Co 5

SINTERED MAGNETS

Ce (Co,Fe,Cu)>5 S m,C • ( Fe, C u ) 6 - 7

V//////////////////A

(Sm,HRE)(Co, Fe,Cu,'I'M ) 7 -8. 5 V/'///////////////'A (Sin, CeXC°, Fe,Cu,TM)6.5-8.5

V////////////////////////////~I---_-_]

Nd~>Fel4B ~ / / / / / / / / / / / / / / / / / , / ~ / / d AND MODIFICATIONS

o

-----_----- -- ------

-- - -_".

I

I

f

I

I

I

I

~

I

I

~

,o

,s

20

2~

30

3~

40

45

~o

[MGOe] Fig. 23. Ranges of static energy product values for important types of rare-earth magnets. (After Strnat 1986.) Shaded bars, commercial products; dashed extension, laboratory records.

R-Co PERMANENT MAGNETS

173

6O [ppm] 5O

4O ILl 0 Z Z rn

I

30

20

I0 ,

Y LO 39 57

Ce Pr 58 59

Nd Pm Srn Eu GO Yb 60 61 62 6.3 64 65

Dy Ho Er Tm Yb 66 67 68 69 70

Lu 71

RARE- EARTH ELEMENT Fig. 24. Natural abundance of the rare-earth elements in the earth's crust. (After Greinacher and Reinhardt 1982.)

An overview of all these alloy types and the range of properties available from each is given in fig. 23. Here we use the common figure of merit for magnet performance in static devices, the energy product (BH)max , to compare relative "magnetic quality". This is, of course, a superficial comparison, and the ranking of these magnets for specific application requirements using other criteria may be quite different. In each case, the shaded bar indicates the approximate range of production values, while the dashed extensions show laboratory record values reported for some of the magnet types. Data for polymer-bonded products are also included, and so are, for comparison, values for the new REPM based on N d - F e - B . To put into better perspective some of the comments made above concerning the relative availability or scarcity of the different rare earths, fig. 24 presents an estimate of the abundance of all the lanthanides and yttrium in the Earth's crust (Greinacher and Reinhardt 1982). While the absolute values given are very questionable, the relationship between the several elements appears to be fairly realistic. (Although the balance can also differ significantly between major ore deposits.) The overall trend is that the abundance falls off with increasing atomic number, with a superimposed pattern that the odd-numbered elements are much less abundant than their even-numbered neighbors. Cerium is by far the most abundant rare earth, with Y, La and Nd each being about half as abundant as Ce and thus in fairly good supply. Samarium is an order of magnitude scarcer than Ce on the average. (It is, in fact, by another factor 3 to 10 scarcer in the main

174

K.J. STRNAT

bastnaesite ore bodies which yield most of the present rare-earth production.) Since the rare earth elements always occur together in nature (albeit in two or three major ore types of different compositions) and must therefore be processed together, it is obviously more economical to use mixtures, especially the Ce-mischmetal, instead of the individual elements, where this is possible.

3. Magnet manufacturing technology 3.1. Magnet fabrication methods - A n overview

We shall now briefly review a number of techniques that have been used in either factory or laboratory to prepare permanent magnets of the various suitable R E - T M compositions. They generally require that one or more alloys be prepared first in a separate step. It is therefore convenient to discuss alloy preparation and magnet fabrication methods separately and sequentially. Alloy preparation from the elemental metals by fusion must be done in an inert gas atmosphere (argon, helium) or in vacuum. (During vacuum melting of S m - C o alloys, some of the Sm is inevitably lost by evaporation.) Melting in air or under a slag is not feasible because of the high reactivity of the rare earths. Small-scale laboratory methods used are cold-crucible arc melting, cold-boat induction and levitation melting, fusion in resistance-heated crucibles, and introduction melting in alumina crucibles typically followed by chill casting. Only the latter method is used in commercial production. Near-eutectic RE-Co(Fe) master alloys with a low melting point can be produced by a fused-salt electrolysis method called "electrowinning". These must be remelted into useful magnet alloys with additions of Co and other transition metals, as required. An economical and commercially employed alternative to melting are the calciothermic methods of preparing the alloy in the form of a spongy powder from a mixture of the oxide and elemental metal powders. Herget (1975) and Jones (1987) wrote excellent overviews of all these methods. The production of magnets from the alloys on an industrial scale is generally accomplished by the powder metallurgical method of pulverizing the alloys to a particle size of a few microns, subsequent compaction, sintering and a post-sintering heat-treatment. Alternatively, bonded magnets are made by crushing either the cast alloy or a pre-sintered body to a somewhat larger particle size, then consolidating the powder into a magnet by mixing it with one of several polymeric binders and either die pressing, or injection molding, extruding or calendering the mixture into the desired shape. These methods have been discussed in detail by many authors. [For examples of overviews see Bohlmann (1987), Ormerod (1985), and Strnat (1978c).] Many other magnet preparation methods have been used in research laboratories. Some were developed to a considerable degree of sophistication but are not now used in commercial magnet production, although a few may have practical potential for special application areas. [Cold-compaction of SmC% powder to high density produced the first 20 MGOe magnets, but their properties were not stable enough (Westendorp and Buschow 1969).] Hot-pressing as an alternative to cold-compac-

R-Co PERMANENT MAGNETS

175

tion and sintering could certainly become a commercial process for REPM production as it has in other areas of powder metallurgy. It has also been tried as a method of bending pre-sintered SmC% in a ductile state at high temperatures (Doser and Beach 1974). Hot isostatic pressing (HIP) has been used either as the primary compaction method or as an additional densification step after sintering. Perceived improvements (higher Ho and better elevated-temperature stability due to finer grain and the complete absence of voids) do not appear to justify the higher processing cost. Casting magnets from bulk hardenable alloys, in analogy to the casting of Alnico, was indeed tried on a small commercial scale around 1970 (Nesbitt and Wernick 1973), but it did not succeed because too many parameters need to be controlled simultaneously. Slow directional solidification has produced Cu-hardened magnet samples of respectable properties and was developed as a potential method for making mechanically stronger magnets that contain ductile metal dendrites (Glardon and Kurz 1978). But this, too, has remained a laboratory curiosity. Plasma spraying has been used to produce very fine, grained magnets from several RCo 5 alloys that can have record high coercive force and are also said to have low magnetic losses at high temperatures (see, e.g., Kumar et al. 1978). Sputtered films with useful coercivities and moderately high energy products have been prepared from Sm-Co and several other R E - T M alloys (Cadieu et al. 1985). There are some special microwave device applications where such thin magnets will be useful, perhaps as part of integrated circuits in conjunction with semiconductors and conductive films. Finally, rapid quenching methods such as melt spinning (Lee et al. 1985), and the so-called liquid-dynamic compaction of an atomized molten alloy (Tanigawa et al. 1986), techniques recently developed for the production of N d - F e - B alloys in very fine-grained form, might also have some promise for the preparation of very stable, high-Hc R E - C o magnets. The main problem with rapid solidification and several of the other techniques mentioned above is that they yield magnets with more or less isotropic properties. While this may be desirable in some applications, the high cost of Sm-Co-based alloys generally dictates that they be used in a grain-oriented form to take the best economic advantage of their high energy-product capability. Thus, magnets having a crystal texture with a magnetically preferred axis or plane must remain a development objective whose attainment could make some of these latter laboratory methods technologically useful. This might be done in the initial alloy formation step as, e.g., by applying an orienting magnetic field during film deposition on a cool substrate. Or it may be possible to produce a useful deformation texture, without the use of a magnetic field, by adapting techniques such as the "die upsetting" recently developed for rapidly quenched N d - F e - B (Lee et al. 1985) or hot-extrusion as used with M n - A 1 - C (Sakamoto et al. 1979).

3.2. Commercial alloy production Two methods are now used on an industrial scale, induction melting and calciothermic co-reduction. Procedural variations exist within each (Jones 1987, Herget 1987).

176

K.J. STRNAT

[ "-7 I REDUCE REz03 I I DISTILL SAMARIUM I

RE - METALS - TM (SM) (Co)

PART OF CHARGE MAY ALSO BE MASTER ALLOYS (NEAR-EUTECTIC RE-Co, FERROZlRCONIUM, ETC,)

; I INDUCTION MELT [ I CAST INGOT

]

HOMOGENIZE

PREFERABLY IN WATER-

]

BELOW PERITECTIC TEMP, 1150-1200°C (SMCo5) II00-I150°C (MMCo5)

]

JAW CRUSHER, ETC,, To SEVERAL MM

I CRUSH

1350-1450°C (SM-Co) 1200-1400"C (MM-Co) IN VACUUM OR AR-HE COOLED STEEL CRUCIBLE

I I

METALLOTHERMIC - SM WITH MM OR LA, OTHER RE WITH CA, ELECTROLYSIS - RE OR RE+TM

I m

COARSE GRIND

I I

FINE GRIND

/J

Disc, CONE, ROLL MILL, ETC,, To 50-200 pM STORE IN DRY AR OR N2 BALL, ATTRITOR, JET MILL, STORE IN DRY AR OR N2

Fig. 25. Process outline for the production of rare earth-transition metal magnet alloys by induction melting.

Figure 25 shows the general process steps for the melting route. All components of the alloys must be in metallic form, either as elements, or as master alloys which may be available more economically. Examples of the latter are R E - T M eutectic alloys (except with Sin) prepared by electrowinning, and Fe-Zr or F e - T i which are standard products for the steel industry. The RE metals used are made either metallothermically by reducing RE-oxides with calcium (and in the case of Sm with La or mischmetal) as the reductant, or by molten-chloride electrolysis. Electrolytic methods do not work with samarium because of its stable divalent state. Samarium is usually further refined by vacuum distillation, which is easy because of the low boiling point. The charge is induction melted in an oxide crucible (typically A1203) at temperatures between 1200 and 1450 ° C. The most volatile and reactive components, such as the Sm, are added last to minimize their time at high temperature. There is always some loss of the RE component from the metallic part of the ingot (and of Zr, if present) due to reaction with the crucible, oxide or carbide formation in the melt, and by evaporation of Sin. Corrections must be made by adding excess RE and Zr based on experience. (Picking up 0.1 wt% oxygen corresponds to the removal of about 0.63% Sm from the alloy.) Vacuum induction melting permits very low oxygen contents in the ingot, typically < 500 ppm. The melt should be chill cast to

R - C o PERMANENT MAGNETS

R/D PROCESS

(OR)

RE203(SM2O3) TM (COBALT) CALCIUM OR OR CAH2

177

K0R PROCESS RE203 (SM203) OXIDES AND TRANSITION TM OXIDES (Co304) METALS - FINE POWDERS TM (COBALT) CALCIUM METALCALCIUM GRANULATE

[ BLEND (PELLETIZE) I

PELLETS

FORKOR

I ] REACT IN FURNACE l

i

R/D-115O°C/3H, IN H2 KOR-1000°C/3H, VACUUM

I

HYDRATE CA0, CA ]

WET N2 AT ROOM TEMP,

i

WASH OUT CA(0H)2 ] i I ACID RINSE I

WATER IN AIR AT RT DILUTE (ACETIC) ACID

I

FINAL RINSE I

]

WATER AND ALCOHOL

I

VACUUMDRY

I

RTT° 5O°C' STORE IN DRY AR OR N2

[~

FINE GRIND -7

I

J

BALL, ATTRITOR, JET MILL, STORE IN DRY AR OR N2

Fig. 26. Process outline for the production of rare earth-transition metal magnet alloy powders by calciothermic reduction and solid-state diffusion. R / D is reduction-diffusion, and KOR is Koreduktion or co-reduction.

prevent segregation and to keep the grain size small. In preparing "base metal" alloys with compositions near stoichiometry, say SmCo 5, a homogenization annealing step just below the 1-5 peritectic is sometimes used. When 2-17 alloys are made for bulk hardening and subsequent use in bonded magnets, a solutionizing step in the same temperature range, around 1150 ° C, is required. "Sintering aid", typically a S m - C o alloy with 40-60 wt% Sin, is not homogenized. The alloys, which are quite brittle, are then crushed and milled to a coarse powder for storage in a dry protective gas. When the alloys are prepared at one plant and magnets produced at another, they are usually shipped in this coarse powder form; however, some vendors sell " r e a d y to press" powders with particle sizes in the 3-7 ~m range. There is some danger of fires as well as deterioration of quality associated with this practice. Figure 26 is a flow chart for making alloys by calciothermic co-reduction. This method bypasses several steps in the older a l l o y / m a g n e t production sequence. It is said to bring significant cost reductions by introducing the RE component as an oxide and yielding the alloy in the form of a powder ready for the final milling step.

178

K.J. STRNAT

In the " r e d u c t i o n / d i f f u s i o n " version of the method, first developed by General Electric in the USA (Cech 1974), all the cobalt and any other T M is introduced as a fine metallic powder, blended with calcium granules, and the mixture reacted under a hydrogen atmosphere. The reaction according to the equation ri R203 + 10Co + 3Ca 115°°c' 3hours 2)2RCO5 + 3CaO,

consists of a reduction of the oxide to the rare earth metal, R, and a concurrent solid-state diffusion reaction that forms the magnet alloy, RCo 5 in our example. After cooling to room temperature, the CaO (and any excess Ca) is converted to calcium hydroxide in wet nitrogen gas and washed out by a sequence of rinses with water, a dilute acid, water again, and finally alcohol. This is possible because the alloy forms as an open-pore crystalline sponge. The friable alloy cake is then crumbled, vacuum dried, and is now ready for the final milling step to micron-size particles. Figure 26 also shows another, similar process developed by the Th. Goldschmidt A G in Germany (Domazer 1974). This "co-reduction" or " K O R " method introduces a portion of the 3d-transition element (Co, etc.) also as an oxide, the rest as metal powder. This provides more internal heat in the reacting mass. Also, the reaction vessel is evacuated and the calcium vaporized. No hydrogen is involved. The reaction is represented by the following equation: RaO 3 + nCo304 + (10 - 3n)Co + (4n + 3)Ca 1°°°c' vac'-)2RCo5 + (4n + 3)CaO 3 hours

(n = 0,

, ~-) ....

The further processing of the spongy alloy cake is basically the same as for the R / D alloy. Both methods are, in principle, suited to the production of any R E P M alloy, but the adaptation to Sm(Co, Fe, Cu, Zr)x and now N d - F e - B alloys required extensive development efforts. The R / D or K O R powders have much higher oxygen contents than induction melted alloys, about 1000-2500 ppm by weight. This requires compensation by adding extra RE (Sm), and there are more oxide particles in the final magnet, slightly reducing its remanence.

3.3. Manufacture of sintered magnets 3.3.1. Introduction If the highest remanence and energy density possible for a given alloy composition and good property stability at elevated temperature are desired, the magnets must be nearly 100% dense and have a high degree of grain orientation. These requirements have indeed been present in most REPM applications to date, particularly for the supply-limited and costly S m - C o compositions. They can best be achieved by sintering powder compacts, and so this is the principal manufacturing process today for all 1-5 and 2-17 type magnets. We shall first discuss the typical sequence of steps used for SmCo 5 and similar magnets, then the modifications needed in making bulk-hardened " 2 - 1 7 " .

R-Co PERMANENT MAGNETS

ALLOYPREPARATIO 4N I

I

CRUSHING POWDERMILLING

179

MELTING/CASTING, OR Co-REDUCTION WITH CA

BALLOR JETMILLING To 3-10UM RANGE

I

POWDERBLENDING I

I

MAG,ALIGNMENT ] COMPACTION

COMPOSITION ADJUSTMENT BY 2(3)-ALLOY PROCESS

PREMAGNETIZING, THEN DIE- OR [SOSTATIC PRESSING

I

I

SINTERING

I

I I HEATTREATMENT I

I [MACHINING

(SEE FOLLOWING FIGURES FOR DETAILS)

]

ABRASIVE SLICING, DICING, GRINDING, EDM-CuTTING

]

PULSE CHARGING BEFORE OR AFTER DEVICE ASSEMBLY

I

[MAGNETIZING

IN VACUUM, AR, HE, OR

H2 ATMOSPHERE

Fig. 27. Processoutline for the production of rare earth-cobalt magnets by the sintering route. A generalized process outline is shown in fig. 27. Alloy preparation has been discussed and we shall assume that the alloys are available as coarse powders of several hundreds of microns particle size. To achieve the desired composition, several alternatives are available. In the single-alloy method the starting alloy must be such that after all chemical shifts that occur during melting or calciothermic reduction, and during later steps by oxidation of the powder, the composition of the finished magnet is kept within narrow limits. This requires careful analysis and very close process control throughout the production sequence, and it is difficult and expensive. Therefore, it has become common practice to work with at least two alloys that are milled separately and blended to make compositional adjustments as required by fluctuations in the raw materials or the production cycle. These should be based on frequent in-process chemical and magnetic analysis. For " R C o s " magnets these alloys are a "base metal" near the 1-5 stoichiometry, and a "sintering aid" that is now usually a binary S m - C o alloy n e a r the Sm;Co 7 composition (but with no attempt made to produce a single-phase compound of that ratio). The latter serves two purposes: it allows easy composition control, especially adjustment of the Sm level for oxidation/vaporization losses, and it provides a small amount of a phase that temporarily melts during the sintering step. Such "liquid-phase sintering" accelerates the shrinkage and promotes the required high

180

K.J. STRNAT

ultimate density of the finished magnet body. For more complex magnet compositions, e.g. those of temperature-compensated magnets, a third alloy powder component may be used as a convenient way of adjusting the level of Gd or of a similar additive.

3.3.2. Powder milling and blending The coarse-grained starting materials must be reduced to average particle sizes between about 3 and 7 ~m, depending on the base alloy and often different for the several powders to be blended. This is typically done by either ball-milling in a water-free protective liquid such as toluene, hexane, alcohol or freon, and under a protective gas (nitrogen or argon); or by jet-milling in a stream of dry N 2 o r Ar gas. The objective is to produce single-crystal particles of a narrow size distribution that can be well aligned and will yield magnets of uniform magnetic properties. A variety of ball-malls are in use: rotating-jar-type, attritor and vibratory mills. The powder produced in these must be carefully dried by prolonged slight heating in a vacuum or Ar atmosphere. Residues of the organic liquid introduce carbon into the alloy which is undesirable since it will preferentially react with minor constituents like the Sm or Zr, Ti or Hf in the 2-17 alloys. Jet-milling offers several advantages: the drying step is eliminated and it is easier to achieve uniform particle size. In the 2- or 3-alloy method, the alloys are usually separately milled to their respective optimal particle sizes. They are then mixed in dry form using tumble blenders. Mixing coarse, individually pre-ground powders followed by a final ball-milling of the blend is also done. As an alternative to purely mechanical milling methods, a chemical technique called hydrogen decrepitation has also been used for the comminution of SmCo5 and 2-17 alloys (Harris et al. 1979, Kianvash and Harris 1985). This makes use of the phenomenon that RE-TM crystal lattices can absorb large amounts of hydrogen, swelling and breaking up in the process. Repeated absorption/desorption cycles can produce fine powder, and H-decrepitation has also been combined with mechanical grinding in the hydrided condition.

3.3.3. Magnetic alignment and compaction The most useful magnets are anisotropic with a single preferred axis of magnetization, requiring that the c-axes of all grains in the sintered magnet are aligned more or less parallel. This is achieved by magnetizing the single-crystal particles of the base-alloy powder, applying a magnetic field strong enough to reorient the particles against frictional and magnetic interaction forces, and then compacting to immobilize them in this state. The crystal texture is maintained and usually even improved through the subsequent sintering and homogenizing if properly conducted. The powder compaction is done either by die-pressing or in an isostatic press, sometimes by both in sequence. Dies are used in which the field is either parallel or perpendicular to the pressing force, and occasionally a more complex aligning-field geometry is needed. Perpendicular pressing usually results in better particle orientation than parallel pressing, and isopressing is best. However, particle shape and size distributions, the degree of pre-magnetization, the level of the orienting field, compaction pressure, as well as the external size and shape of the compact, all

R-Co PERMANENT MAGNETS

181

influence the degree of alignment. The pressing force must be high enough to give the "green" body sufficient integrity for handling; but too high a pressure will cause misorientation of particles (except when it is hydrostatic pressure). The orienting field is usually a DC field of 7-12 kOe, sometimes augmented by a superimposed field pulse. In isostatic pressing, only a pulsed field is used, which can economically be made 3-5 times larger than the typical steady holding fields in uniaxial die pressing. An interesting dynamic compaction method that combines pulsed-field orientation and impact pressing can significantly improve the degree of alignment in parallel pressing (Tawara et al. 1985). When isotropic magnetic properties are sufficient or desired, no pre-magnetizing or orienting field is used. However, it is almost impossible to prepare a truly isotropic magnet by die pressing, since there are always some minor mechanical alignment effects. On the other hand, no way has been found as yet to prepare R E - T M alloy particles with a clearly developed platelet or needle shape related to the crystal symmetry in such a way that it would allow purely mechanical orientation, resulting in a magnetic preference direction. (This is possible in the fabrication of certain bonded hexaferrite magnets.) 3.3.4. Sintering and heat-treatment The sintering step can be carried out in a high vacuum, in a dry atmosphere of Ar of He (free of O 2 and N 2), or in hydrogen gas. Sometimes these are used sequentially. Hydrogen can improve and accelerate the densification. However, it must be removed from the sintering furnace before cooling because of the tendency of R E - T M compounds to absorb it and swell, which can physically destroy the magnet. Even small amounts of hydrogen retained in the lattice will adversely affect the magnetic coercivity (Zijlstra and Westendorp 1969). The objective of sintering is to produce a homogeneous body of high density while controlling crystal growth and texture. In SmC%, and in other magnets where grain-boundary pinning or "nucleation" determine He, appreciable grain growth must be prevented. The compromise between these conflicting requirements demands close control over sintering temperature, T~, and sintering time. For SmCo 5 the optimum T~ = 1130-1150 ° C, for other RCo 5 it is lower, down to about 1050 ° C for high cerium or copper contents. In the " 2 - 1 7 " and other magnet alloys that are magnetically precipitation hardened, grain growth is a secondary consideration. Typical sintering temperatures range from 1190 to 1220 ° C for Sm-based magnets, and again lower when the Ce content is high. Sintering times of about 30-90 rain are used. A typical sintering cycle for SmCo 5 is shown in fig. 28. Immediately after sintering, the coercivity and demagnetization curve shape are poor. Annealing above the eutectoid temperature, followed by rapid cooling through the critical range where decomposition occurs fairly quickly (from about 800 down to 500 ° C) will develop the best permanent magnet properties. The optimum temperature for this post-sintering heat-treatment is 850-900 ° C, as can be seen from the variation of MHc with the temperature of such an anneal, shown in fig. 29. The demagnetization curve shape is quite sensitive to the rate of cooling after this anneal, see fig. 30. A

182

K.J. STRNAT

1200

SINTER ANNEAL

1000

G o v

w Q: 800

600 LU 0-

400 I-"

200

FURNACE COOL

;y

QUENCH I

I

5

10

I

TIME (HOURS)15

20

Fig. 28. Typical temperature-time profile for the sintering of SmCo5 magnets. (After Paladino et al. 1975.)

systematic s t u d y of the influence of the cooling rate was r e p o r t e d b y den B r o e d e r et al. (1974); their results are shown in fig. 31. It c a n also b e seen t h a t the p r e s e n c e of excess Sm in the m a g n e t can s t r o n g l y e n h a n c e the coercivity. F o r n e a r - s t o i c h i o m e t ric 1 - 5 c o m p o u n d s fairly r a p i d cooling is essential. Oil a n d w a t e r q u e n c h i n g have b e e n used, b u t t h e y cause cracking p r o b l e m s . A " g a s q u e n c h " or even c o o l i n g in s t a g n a n t gas o u t s i d e the furnace are generally sufficient.

I

r

i

i

~

i

i

i

i

I

L

i

i

MHc (kOe) 25

2O

15

10

5 0

i

700

750

800 850 900 1000 1000 ANNEALING TEMPERATURE (°C)

i

1050

Fig. 29. Dependence of the intrinsic coercive force of sintered and quenched SmCo5 magnets on the temperature of the post-sintering anneal. (After Paladino et al. 1975.)

R - C o PERMANENT MAGNETS

183

4 ~ M (kG) 10

6

COOL

/

I

20

15

4

COOL

I

I

-H (kOe)

2

[

10

0

5

0

Fig. 30. Demagnetization curves of sintered SmCo 5 magnets cooled from the annealing temperature at different rates. (After Paladino et al. 1975.)

The Cu-hardened magnet types require a different, more complicated heat-treatment to develop the precipitate structure. For R(Co, Fe, Cu)5 no rapid quenching step is needed, an advantage in making larger shapes that are difficult to cool rapidly throughout their volume. For the "2-17" magnets, however, there is again a rapid cooling step required, after the homogenization treatment, and the heating cycle is generally more complicated and prolonged. This adds significantly to the production cost since higher-temperature sintering furnaces and altogether more furnace capacity with precise temperature control is needed. 7O

MHc

tOe] 5O

J

SmCo4.2 f

J

6j io f

~SmCo4.7 30

o~_ ~ . . _ . ~ L (0)

0.1

k_______ I

I0

I00

I000

COOLING RATE [°C/s] Fig. 31. Dependence of MITeon the average cooling rate through the range from 800 down to 500 ° C for sintered magnets SmCo 5 x with different Sm excess over the 1-5 ratio. (After den Broeder et al. 1974.)

184

K.J. STRNAT T

[-c]

1220 °

1200

i 1 ~

I

8OO

°

I

850"

~ ; ~ L O W COOL - I 0 HRS.

4OO

~ 3 - 5 HRS. ~" • I

-400* - -0 - 1 0 HRS.

I

QUENCH

0-

1

• "OUTGAS

,~

1

TIME

'

I

/, I

J7

I

FkGl I I

/.E o j

_-JJ

Fig. 32. Typical temperature profile for the sintering and heat-treating of "2-17"-type Sm(Co, Fe, Cu, Zr) 7.2_s.5 magnets.

Figure 32 shows a generic temperature-time cycle for the production of sintered magnets of compositions Sm(Co, Fe, Cu, Zr) 7.2- 8.5. The sintering must be followed by a solutionizing anneal at a temperature in the 1130-1175 ° C range, the precise level depending on the exact composition and being quite critical. This produces the single-phase, 2 - 1 7 R solid solution from which the cellular microstructure can then develop during an isothermal aging anneal of several hours near 800 o C. Finally, it is necessary to cool the magnet in a very slow and controlled manner to about 400 ° C (sometimes in many small steps), and it is desirable, but not necessary, to hold it there for several more hours. The hysteresis loops drawn below the time-temperature curve indicate the magnetic properties at different stations: the initial loop indicates poor crystal alignment; it becomes taller, with horizontal tails, indicating improved texture after the sintering and solutionizing, but H c is now very low; the curve remains similar even after the isothermal again near 800 ° C. High coercivity and loop squareness develop only during the slow cooling, and they are slightly improved by the further heating at about 400 o C.

3.3.5. Machining, handling and magnetizing The sintered magnets are hard and brittle. They can only be machined by the techniques developed for Alnicos, ceramic magnets or semiconductors. They can be sliced with diamond impregnated blades, surface ground with carborundum or diamond grinding wheels, or cut on electric spark-erosion machines (EDM). This must be done in a demagnetized state. The high intrinsic coercive force and near-unity relative recoil permeability of most REPM permits their manipulation in the magnetized state without the need for "keepers". The magnets can be pre-charged in open circuit and will not self-demagnetize. However, machining, shipping and handling in the magnetized

R-Co PERMANENT MAGNETS

185

% ] OF 8 0 kOe VALUE I00

60

A'Hc H ¢/ _/ /- /

4o 20

f f "~ 8r after field

demognetizo--

lion 0

5

I0

15

.[kOe]

20

25

Fig. 33. Magnetizing, or "charging" behavior of a typical sintered SmCo5 magnet. (After Bohlmann 1976.)

state are difficult, even nearly impossible for larger high-energy magnets. It is therefore usually necessary to magnetize after partial assembly of a device or machine. This can also present difficult problems. Some of the lower-Hc bulk hardened magnet compositions require only 8-12 kOe to fully magnetize them; so they may be handled like ferrites. However, SmCo 5 and, especially, the high-Ho versions of 2-17 need much higher charging fields, up to 50 kOe (4 M A / m ) . Producing these, even as millisecond pulses, requires novel and expensive magnetizing equipment, and it may even be impossible to achieve sufficiently high fields in a larger assembly. Figure 33 illustrates the magnetizing behavior of a typical commercial sintered SmC% magnet. Magnetizing from the virgin state, it takes about 15 kOe to fully develop the flux density, Bd, at the operating point for maximum energy product. However, developing the best demagnetization curve squareness in the region just before the ~/H~ point requires much higher fields. A figure of merit used for devices operating there, the "intrinsic energy product" (Bill)max is only at 90% of its optimum value after a 25 kOe charging field. For certain microwave tube and actuator types that operate out to points near the "knee" of the intrinsic curve, it is important to magnetize in fields higher than this. And, after the magnet was once charged and then demagnetized in a field (rather than by heating to near Tc), it becomes still very much harder to magnetize, as is indicated by the dashed curve of B r versus H. It is also important to know that incompletely magnetized or partially demagnetized SmCo5 magnets exhibit severe thermal instabilities of the flux, termed "thermal remagnetization" (Livingston and Martin 1983). High-Hc 2-17 magnets typically require over 50 kOe magnetizing field for full development of their remanence and their second-quadrant properties, as is shown in fig. 34. Different from SmC%, this behavior is independent of the previous magnetization history. Also, bulk-hardened magnets do not show the above-mentioned thermally activated flux creeping to any significant degree. The process of magnetizing or re-magnetizing can be assisted by heating to elevated temperatures, where wall pinning becomes less effective and is overcome by smaller applied fields. Figure 35 shows how the B r versus H curve of the same high-Ho 2-17 magnet

186

K.J. STRNAT 2

0 2.4 .~,~

HMAx ,[MA/m]

6

8

,.c 12 Br

MHc

IT] [kS]

/////

1.6 2O ///

0.8- 8

/-

MHc

0.8

0.4- 4

o

0

25

50 HMAX [kOe]

75

I00

0

Fig. 34. Dependence of B r and MHc of the magnetizing field for high-coercivity sintered "2-17" magnet of composition Sm(Co, Fe, Cu, Zr)v.5. (After Ervens 1982b.) changes on heating. The magnetizing field needed to develop the full remanence is reduced f r o m over 50 kOe at r o o m temperature to less than 25 kOe at 400 o C.

3.4. Bonded magnets 3.4.1. State of development Bonded, or matrix magnets utilizing R E - T M alloys as their magnetic constituent were recognized as a potentially very useful p r o d u c t family in the late 1960s (Strnat 1967). However, for various reasons, technological as well as economical, their development proceeded m u c h more slowly than that of the sintered magnets. There were m a n y difficulties to overcome, relating primarily to the high chemical reactivity of the alloys in powder form, and inexpensive m a n u f a c t u r i n g methods h a d to be

B r 1.2

H [MA/m] 4 6 I I

2 I

0

8 I

12 B r ,0 [kG]

IT] ,.0

0.8

I

I

/ / / / y ~

0.6 0.4

/// J

/ MAGNETIZED AT: -

8

I - ROOM TEMP. 2-200°C 3- 4oooc

4

0.2 0

6

[

2

0

I 25

I I 50 75 H[ kOe]

I IOO

0

Fig. 35. The influence of temperature on the magnetizing behavior of the 2-17 magnet of fig. 34. (After Ervens 1982b.)

R-Co PERMANENT MAGNETS

187

developed taking these problems into account. Early products based on fine particles of SmCo 5 (2-10 ~m), and bonded with epoxy resins (Taylor and Wainwright 1976) or with thermoplastic resins such as polyethylene chloride or ethylene vinyl acetate (Kamino and Yamane i976, Suzuki et al. 1978), deteriorated quickly above 50 °C and were magnetically not very stable even at room temperature (Mildrum and Wong 1976). Various methods for coating the particles with metallic or organic diffusion barriers, careful selection of compatible polymeric binders, etc., have brought much improvement since then, and useful SmC%-based bonded magnets are now on the market (Suzuki et al. 1979, Satoh et al. 1985). However, the most important breakthrough was the introduction of alloys bulkhardened with Cu. If used as coarser powders ( - 20-100 ~tm) they do not suffer from the severe sensitivity of H c to the particle surface condition, and so they are much more stable (Strnat et al. 1976a,b, Shimoda et al. 1979). The use of a soft metal such as Sn-Pb solder as the matrix is another way of strongly improving the elevated-temperature stability (R.M.W. Strnat and Liu 1982). Cu-hardened alloys close to the 2-17 stoichiometry and with higher Fe content have been developed for bonded magnets (Shimoda et al. 1979, Ray 1984), as well as ways to produce ingots with a uniform columnar grain structure from them, removing the need for using pre-sintered blocks. These developments, together with the recent adaptation to plastic matrix REPM of inexpensive molding methods such as extrusion and injection molding in an orienting field (Hamano 1987, Satoh et al. 1985, Shimoda et al. 1985), appear to have cleared the way for a large-scale indastrial application of the bonded REPM. The current situation was reviewed by a major contributor, Shimoda (1987). 3.4.2. Fabrication methods

There are many possible alternatives for the processing of polymer- or metal-bonded REPM. Figure 36 is a generalized and simplified flow chart of steps common to different techniques. However, it pertains primarily to the magnet shaping by die-forming, rather than the injection molding or extrusion methods. In the upper half of the chart, the left branch refers to the utilization of reground pre-sintered blocks or of scrap from the sintered REPM production. The right branch shows the preparation of powders from coarse-grained cast ingots of bulk hardened alloy. In the lower half, the left branch represents the common method of preparing first an alloy-binder mixture and then forming it into the desired magnet shape. The binder can, in principle, be a thermoplastic, rubber, thermally or chemically curable polymer, or a low-melting soft metal. The magnets made by this method range from isotropic to imperfectly anisotropic and have relatively low to moderate values of energy product. The procedure in the right branch of the chart is capable of yielding the densest and best-aligned magnets with the highest possible energy products, up to 20 MGOe in laboratory, 17 MGOe in production. This was accomplished by using a high-remanence 2-17 alloy, paying careful attention to the particle size distribution, orienting it in a high transverse field, avoiding misorientation during die compaction, and then infiltrating a liquid epoxy resin into the pores of the compact (Shimoda et al. 1980, Shimoda 1987).

188

K.J. STRNAT

ALLOY PREPARATION (MELTING AND CASTING) 1

HEAT TREATMENT FOR SOLUTIONIZING AND GRAIN GROWTH

I MILLING TO ~M, FIELD-COMPACTION (DIE- OR ISOPRESS)

I

I

HEAT-TREAT INGOT FOR MAG, HARDENING

SINTERIi~G MAG, HARDENING

1

I

COARSE GRINDING]

I

1

I [ PARTICLESIZING1 I

I

I POWDER BLENDING ADD SOLID BINDER

I

I POWDER BLENDING

[

COMPACTION WITH OR WITHOUT FIELD

COMPACTION OR FORMING WITH OR WITHOUT FIELD I

LIQUID BINDER IMPREGNATION

CURING OR MELTING OF BINDER ]

i (MACHINING, FIliAL MAGNETIZING) Fig. 36. Generic outline of process steps used in the production of bonded REPM.

3. 4. 3. Subtypes and their properties The best compression-molded bonded 2-17 magnets available today offer roomtemperature properties equivalent to medium-grade sintered SmCo 5 and are far superior to any older, non-rare earth magnet type. Injection-molded magnets naturally have significantly poorer demagnetization curves; their commercial attraction is the low cost of processing. Plastic bonded R E P M are now produced with energy products covering the range from about 3 to over 17 MGOe, from 2-17 alloys on the upper end, through SmCo 5 in the middle, to alloys with some mischmetal substitution for Sm in the low property regime. Figure 37 shows the range of B, H-demagnetization curves available and compares them with the characteristics of older types of bonded magnets. While metal-matrix R E P M are not yet commercially produced, it should be possible to duplicate with them the same range of room-temperature magnetic properties, while improving the high-temperature stability and mechanical strength and increasing the thermal conductivity to fit especially demanding applications. They are, of course, electrical conductors, like

R-Co PERMANENT MAGNETS -H [ k A / r n ]

400

189

200

0.8

B

B

[T]

[kG]

0.6

0.4

Q2

2

o

3 -H [ k O e ]

CHIPPED ALNICO

Fig. 37. Comparison of the second-quadrant demagnetizationcurves of commerciallyproduced bonded REPM and older types of matrix magnets. the sintered or even the high-density polymer matrix magnets. (Low-density plastic magnets with a high binder content can be made to have high electrical resistance, which is occasionally required.) Metal-matrix magnets have also been prepared with a low packing fraction of the 2-17 magnet alloy, a S n - P b binder and a pure Cu filler, by die-compaction without heat, to obtain magnet sets with precisely controlled variable saturation (R.M.W. Strnat et al. 1987a,b).

3.4.4. Molding, machining, handling and magnetizing The most important advantage of bonded magnets is the possibifity of molding them to the final required shape with close dimensional tolerances. In mass production this means significant savings in processing costs, and since there is virtually no material wasted as cutting scrap, or lost in grinding, such molding methods will also help conserve the expensive and supply-limited magnet alloys. It is further possible to produce in a single molding operation composite parts and

190

K.J. STRNAT

subassemblies, such as rotors or stators for small motors, which incorporate nonmagnetic structural parts or iron of the magnetic circuit. If machining is necessary, however, it can often be done with conventional machine tools such as a drill or lathe. Because of the strongly abrasive nature of the magnet alloys, hard-faced tool bits are required. The necessity for machining may arise from particularly tight tolerances for dimensions or surface smoothness. Surprisingly, though, there are also production methods in use for very small parts which require much slicing, dicing and drilling (Shimoda et al. 1979). With regard to handling, bonded magnets cause much fewer problems in production than the sintered REPM. Depending on the binder content, their mechanical behavior ranges from somewhat ductile (with metal matrix), and elastic or even flexible (with plastic and rubber binders) to moderately brittle when the alloy content is very high. Since bonded magnets also have lower remanence than their sintered counterparts, the forces encountered are lower, and the net effect is that there is much less danger of breakage or injuries in device assembly. On the other hand, anisotropic bonded magnets come off the production line in the magnetized state, and so it is indeed desirable to manipulate them that way to avoid additional demagnetizing and recharging steps.

4. Properties of commercial magnets 4.1. General comments

Our objective in this section is t o summarize the information previously developed in this chapter about magnet types and their characteristic properties. The emphasis here is on those magnets which are now in commercial production or in an advanced state of manufacturing process development. An attempt is made to present systematically, mostly in tabular form, data for the salient properties of interest to design engineers looking for the most suitable permanent magnet to use in a circuit. This must include not only magnetic properties but also mechanical, electrical and thermal characteristics that are important in different applications. This is indeed not an easy task and the result cannot by very satisfactory at this time for several reasons. One reason is simply the bewildering multitude of magnet types which are at least in a small-scale pilot production. T h e discussions in the preceding sections have shown that there is an almost infinite variety of R E - C o based alloy compositions that can be given useful permanent magnet properties. Several different metallurgical microstructure types cause magnetic hardening by several mechanisms, and these lead to quite different magnetization and demagnetization behavior. And there are the two basic manufacturing techniques, sintering and bonding; either - but especially the latter - allows many variations which lead to a wide range of engineering properties and their thermal and temporal stability. A set of properties cannot be simply correlated to an alloy composition alone. For a variety of reasons - special market niches, raw materials supply, other economic considerations, and patent protection - REPM producers worldwide have chosen to manufacture a large number of different magnet types. There has been little

R-Co PERMANENT MAGNETS

191

industry coordination as yet, although considerable progress toward standardization has been made with regard to sintered SmC% and Sm(Co, Fe, Cu, Zr)_7. 5 compositions. Another problem lies in the fact that many magnet manufacturers do not reveal details of the composition and heat-treatment of their products, while indeed any effort to compile design data must in large part rely on commercial product literature as the source. Many properties, especially the nonmagnetic ones, have not yet been reliably measured on well-defined samples. Therefore, some of the data presented in the tables must be taken with a grain of salt. Often only property ranges can be specified; and where only numbers from a single measurement have been found, they are shown with an approximate sign or the problem is explained in a footnote. The data stem from many sources, scientific papers, manufacturers' literature and unpublished results in the author's laboratory. No attempt is made to identify the sources of specific values. However, the foregoing discussions of specific magnet categories identify many useful references for technical data, especially some of the technical review articles such as Ervens (1982a), Nagel (1980), Mildrum et al. (1981) and Strnat (1983). With regard to the temperature variation of magnetic properties, their behavior under thermal cycling and their long-term stability, relatively much information has been generated for elevated temperatures, to limits between 200 and 400 o C. This allows the designer to define fairly well the maximum useful service temperatures of different magnet types under the restraints of a specific application, such as the operating permeance of the magnet and maximum permissible flux change. While it exceeds the scope of this chapter to report such data in great detail, information is presented for a few important sintered material types which are the magnets of choice for elevated-temperature applications. For special applications the permanent magnetic properties at very low temperatures, down to 4.2 K, are desired, but there is as yet almost no engineering information available. Basic science studies of magnetization, anisotropy field and coercivity at cryogenic temperatures show that for SmC% and Sm-Co-based 2-17 magnets all these quantities increase on cooling to near absolute zero. From this one can conclude that such magnets, when charged at room temperature and then cooled, get only "better". If a small gain in flux density and large increases of the intrinsic coercive force are acceptable, these materials are good choices for low-temperature operation. On the other hand, all N d - C o compounds, and PrCo 5 as well, enter an easy-cone anisotropy state on cooling which will decrease the coercivity and distort the demagnetization curve. Magnet compositions with a high N d content, and to a lesser degree those containing much Pr, are not suitable for use in cryogenic devices. Another concern for special new applications, in nuclear reactors and particle-beam devices (such as synchrotron light sources), is the behavior of magnets in a radiation environment. This, too, is essentially outside the scope of this chapter, and no data are included. Suffice it to say that high doses of neutron and ionizing radiation have been reported to have fairly severe degrading effects on the open-circuit flux of 1-5 and 2-17 sintered S m - C o magnets. This is in contrast to the results of old investigations which showed virtually no radiation effects on Alnico and ferrite magnets. Much of

192

K.J. STRNAT

the observed degradation may be an indirect effect of internal heating caused by the absorbed radiation and should thus depend on the size and surface-to-volume ratio of the magnets. If this assumption is correct, the effects may be lessened by effective cooling, provided the application permits this. At the time of this writing, such radiation-effect studies are in progress at several reactor and accelerator laboratories.

4.2. Magnetic properties of sintered magnets Table 3 lists the nominal compositions and salient permanent magnet properties of twelve rare earth-cobalt-based magnets and, for comparison purposes, also the properties of a magnet based on the Nd2Fe14B compound. All formulas represent materials that were developed to near production maturity, and most are in present commercial production. The property data were usually measured on selected samples prepared under laboratory conditions, so they must be taken to represent the limiting properties that are practically possible for a given alloy. Even the best production magnets must be expected to have somewhat poorer performance, perhaps about 10% lower energy products. Figure 38 is a comparison of the second-quadrant B,H-curves for typical sintered production REPM, comparing them with demagnetization curves for major older magnet types. The temperature coefficient of MHo was obtained from measurements of demagnetization curves at several temperatures, each traced after fully remagnetizing the sample at room temperature. As an example, fig. 39 shows such a set of curves for a sintered SmC% magnet of average production quality. The coefficient of B r is taken from the reversible remanence versus temperature curve traced after a first heating cycle to 100°C, i.e., the so-called irreversible loss has been eliminated. Both

[kA/m] 1.4 800

B 1.2

6,uO- 400 200

Alnico5~ 12 B

SmlCo,Fe,Cu,~

[kG]

,o

•

0.8

0.6

)14

8

~

6

0.4

4

0.2

2

012 I0

8

6 4 -H[kOe]

2

00

Fig. 38. Comparison of REPM with older permanent magnets. B versus H demagnetization curves typical of high-grade commercialmagnets of the types indicated.

R-Co PERMANENT MAGNETS

193

r2

O

I O t'N

I

t',l

i.!

O

Q ©

6

r,.) I

r.s.l

e,,

t,-

g

##

194

K.J. STRNAT

(-B/Hor-B/y-gH)

I~

2%2.5\3\

51 I01

[kG]!T]2 12

I0 - 1.0

B (B-H)

O

_60o C

+ 25=C

+ i00o c

+ 200oC

////.

r_ 1 LkOeJ 25 r,. L~, ','" L~*mjA 2 0 0 0

,

0.6

4

0.4

oOo

20

15 15--00

6

'

I 0 '0 0

IO - H '

5 5 0' 0

i

0

Fig. 39. Demagnetization curves of a sintered SmCo 5 production magnet measured at four temperatures. (After Mildrum et al. 1981.) Note: The scale markers on top indicate the end points of permeance fines that may be drawn from the origin. Their intercepts with the B versus H-curve define operating points with the indicated unit permeance values, p = - B d / H d = 1,..., 10.

temperature coefficients are calculated from a linear interpolation of the property values at 20 and 100 o C.

4.3. Magnetic properties of bonded magnets

Table 4 is an attempt to summarize the commercial offering of all types of polymer bonded R E - C o magnets in mid-1988. Since this product line is in a state of rapid development, some of the highest values for B r and (BH)max may have been slightly exceeded by the publication date of this book and may be expected to further improve somewhat. It is difficult to identify the exact compositions of the magnetic alloys used from the commercial product brochures. However, it is safe to assume that most of the products in the table contain a " 2 - 1 7 " type Sin(Co, Fe, Cu, Zr)x alloy, with a minority being based on SmC%. The products are arranged according to the basic magnet-forming method used, with a further subdivision by the geometry of the orienting field applied, if any. Considering the large discrepancy in the mass densities of alloy (> 8 g / c m 3) and polymer ( - 1 g/cm3), the density of the magnet is a rather direct measure of the packing density (fill factor) of the magnetic alloy in the magnet body.

R-Co PERMANENT MAGNETS

195

E

"0

0 e-

o c~

~

d

ddd

~

ggg

E

~ ~

d

ddd

',~

~

t"-I t¢3

oo

~D

~o

cq

O~

0

tr~ uq

E

e. oo

~ i

d

,

,

c~

ddd

o

L) ~'

,,d

o

-

o .u .~ ~.~

~

~oo

~-

f

,,/

<<

.u

~ ~

~

.u

.£

~

~

.&

<

,<

c~

g 0 o

.=~o

~ ~

o o .,~,

.,

~

""

196

K.J. STRNAT

4. 4. Magnetic stability - Thermal cycfing and aging For many applications it is important to know how the magnetic properties change as a function of temperature, and how stable they are over short or long periods of heating to elevated temperatures, especially when exposed to air. This is of critical significance for devices that require that the flux be either constant within narrow limits over the operating temperature range, or at least precisely predictable so that external compensation measures can be taken. For these devices, the flux should also remain invariable over the operating life of the device. The unusually strict stability requirements of microwave tube focusing structures, accelerometers and navigational gyros have motivated a series of careful measurements of the so-called reversible and irreversible flux losses from room temperature up to the point where the changes become intolerably large. On the lower temperature end, several investigations have been extended to - 6 0 ° C. For reports of such data for different REPM materials see, e.g., Mildrum et al. (1974), Mildrum and Wong (1976), Jones and Tokunaga (1976), Bachmann (1977), Narasimhan et al. (1978), R.M.W. Strnat and Liu (1982), Ervens (1982a), Hanna and Walmer (1985), Shimoda et al. (1985), and Li et al. (1988). The measurements needed are different from the previously discussed plotting of major second-quadrant demagnetization curves at several temperatures after full charging of the magnet. The flux changes during short heating-cooling cycles (the so-called reversible and irreversible losses) and the severity of long-term aging effects depend upon the operating point of the magnet on the demagnetization curve, which is usually described by specifying an operating permeance, p = - B d / H d. The irreversible loss at lower temperatures is strictly magnetic in origin and recoverable by remagnetizing, but not simply by cooling to room temperature. It is attributable to thermally activated domain-wall motion under the influence of internal (self-demagnetizing) fields. This loss component increases rapidly with decreasing [ p] (Martin and Benz 1972, Mildrum and Wong 1976). The reversible loss (often also characterized by a temperature coefficient) is almost independent of the value of p, tracking more or less the temperature variation of the spontaneous magnetization. At higher temperatures so-called irrecoverable losses come into play Which are due to slow changes in the metallurgical microstructure and to oxidation, and in bonded magnets also to a softening of the binder and a deterioration of the alloy-binder interface (R.M.W. Strnat and Liu 1982). These effects become significant at 2 5 0 - 3 0 0 ° C for sintered SmCos, at 3 5 0 - 4 0 0 ° C for sintered 2-17, but at much lower temperatures ( 5 0 - 1 8 0 ° C ) for the various types o]~ bonded magnets. Definitions and more detailed discussion of the origin of these losses, with examples of the typical behavior of different older magnet types, can be found in the reviews by Tenzer (1969) and Strnat (1983). Let us now first consider the reversible losses in R E - C o magnets. By way of example, fig. 40 shows for five commercially important types of sintered magnets the reversible flux variations on heating up to 250 ° C, measured at p ~ - 2 . 5 . It is seen that the reversible loss (or the temperature coefficient of Bd, which is the average slope of such curves over a defined temperature interval), is primarily

R-Co PERMANENT MAGNETS 0 -

197

BIH ~ - 2 . 5 ~

G

4-

H \- "- ~

-

H

c 2-17

~ " -~~ / ~,~ LOW-- Hc 2 - 1 7

O _J

~8 _J [a.

Jm 12 \\ ¢m "' Smo.sCeo.5(Co,Cu, Mn)7~\\ \\\ >is J ""

c.oo

0:

(

20 o

I 50

Cu, Fe) 5

I I 1 Ioo 15o 200 TEMPERATURE (=C}

',, I 250

Fig. 40. Reversible temperature variation of the open-circuit remanent flux at an operating permeance of p ~ - 2 . 5 , measured on five commercially important sintered magnet types. Curves traced after thermal prestabilization by heating for 5 h at 250 o C, at 200 o C for Ce(Co, Cu; Fe)s. (After Li et al. 1988.)

determined by the composition of the magnet alloy. Note that the two 2 - 1 7 magnets have similar losses despite their very different coercivity! This flux loss is closely related to the Curie temperature of the main phase, which is of course a function of the alloy composition. Also, the more cerium an alloy contains the lower is its Tc and the greater the reversible loss. Such measurements of flux losses as a function of temperature or time are usually performed on small samples of cylindrical or prismatic shape of a length-to-diameter (or edge length) ratio which is said to correspond to a certain value of p = B d / H d. However, since magnet shapes other than ellipsoids are never uniformly magnetized, and their H d is also not uniform through their volume, it is really impossible to define a single, precise p-value for the usual sample shapes. This is the reason why all p values in the figs. 40-44 are indicated as approximate. The following discussion applies to "irreversible losses" and "aging effects", which are indeed difficult to separate. When magnets are exposed to elevated temperatures for very long periods of time, as they are in the operation of certain devices such as microwave tubes, additional time-dependent losses of the useful operating flux are observed which are called aging losses. The domain-wall creeping responsible for the "irreversible" losses is not only a function of temperature but also of time. At lower temperatures, say up to 100 o or 150 ° C, the magnetic domain structure m a y not reach an equilibrium state for several hours (at the upper limit) or even several hundred hours (near room temperature). Therefore, i f the flux is measured as a function of time at temperature, a gradual drop is observed that is initially fast, then its rate declines, and finally the flux levels off. At that point in time the magnet can be considered "thermally stabilized". But at higher temperatures (with an ill-defined lower threshold that depends on the nature of the magnet)

198

K.J. STRNAT

the effects of metallurgical structure changes such as the eutectoid decompositions of SmCo 5 or the coarsening of precipitates in the Cu-hardened magnets, or of oxidation at grain surfaces, or of binder deterioration, become noticeable. In higher-temperature aging curves they show up as a flux reduction from the earlier stabilized level, and they can eventually lead to a catastrophic loss of properties of the magnet, definitely limiting its useful lifetime. The bulk magnetic property primarily related to this irrecoverable loss is the intrinsic coercive force. Therefore, the flux loss is more severe when the initial MHc is low, and it is worse for lower I Pl values for which the operating point on the demagnetization curve moves closer to MHc. For high-coercivity magnets such as sintered SmCo 5 or S i n - C o - b a s e d 2-17, the time domains of the initial, purely magnetic losses and of the "metallurgical" losses are clearly separated at intermediate aging temperatures, say 150-250 ° C, b y a long plateau. In contrast, for magnets in which structural changes can occur at fairly low temperatures, these regions in the flux versus time curves blend into each other and there may be no plateau at all. This is the case for Ce(Co, Cu, Fe)5 and for many polymer-bonded magnets. Figure 41 shows aging curves measured on the same set of samples as the reversible losses in fig. 40 (Li et al. 1988), at 150 o and 250 o C. The samples were axially magnetized cylinders with L/D ratios corresponding to two rather different permeance values: p ~ - 2 . 5 is characteristic of the operation of magnets in m a n y electric motors, while p = - 0 . 5 is close to the typical operating point of travelingwave-tube focusing magnets. These measurements were made at 25 ° C after the magnets had been heated in air at the temperature and for the exposure time indicated. Note that the initial points, at t = 10 -1 h, thus correspond to a short-term irreversible loss, while the flux reduction at very long exposure times are clearly "aging losses" that must be expected in device operation. However, it is obvious that heating between about 1 and 10 h often stabilizes the flux quite well for a long subsequent time period. Such artificial pre-aging can be employed as a method of thermal stabilization, and aging curves like those shown help determine the time-temperature conditions which will accomplish this (Mildrum et al. 1974). Note also that some authors define an "initial loss" consisting of the reversible plus irreversible losses incurred up to a stabilization period defined for each specific case, typically between 1 and 5 h. The results of another set of very systematic long-term aging studies are presented in the figs. 42-44 (Ervens 1982a). They compare the behavior of sintered SmCo 5 and high-coercivity 2 - 1 7 magnets of large and comparable MH~ values (25 and 24.3 kOe, respectively) at temperatures up to 400 o C, for exposure times to 5000 h (ca. 7 months). These measurements were done at different operating permeances than the curves of fig. 41, namely, p ~ - 1 , near the (Bg)max point, and p ~ - 2 . The ordinate scales show a normalized flux density value for the operating point, with Bd(t)/Ba(O ) = 1 being equivalent to loss = 0 in fig. 41. Figure 42 summarizes the aging behavior for 2-17. This material shows acceptable stability, with < 5-6% loss, up to at least 300 ° C, while it is clearly no longer useful at 400 o C and p = - 1. Figure 43 shows the 1-5 and 2 - 1 7 magnets to be nearly equally stable at 200 o C, while at 3000C the 2 - 1 7 is clearly superior and still useful, while the 1 - 5 ages

R-Co PERMANENT MAGNETS

199 .¢&

I

I

I

I

-~ ~ _ ","

-=

o,

II

I "~. I

~

i

I

T 9

I

I

~

/-~ ~

//-

,= ~

, .o,

' ,/

"'

-2

--

6

o_~.I'i

J

p=.; I

o,

~

_

/

_d~

'~ ~

~i£

/-,..

:,

.~g.o

I

--

~ •

u

-o,

-

I

I>"

~,

?

I

IE

j

\.

I

~

I

_ ,4

I

~

l J_

°it -- ~--

; =

0 9

~"

I

II

o

I

/

,~

o

x~

I.= I ~ ~'~

~/~

,,'>

/ ~

~'. ~

~,II ~,,~-

I

~

<'/

o

,,',

o

E o II

" ~ ~"£

,,',

-

II

~

~,~ o I- ± IIII ~ /:~

~

r/ o~ .

;iIll

T

~

(~

o

.@ .

' I

T~

I/

"

~- I J

/

=,~ ~

ill;',,,

I

I 0

LI..J ~.,

~

I 0

= oC~ "1~

-,::, ~ o =~

o,~ o ~ --:3: l'- ",

~g

To o£ r~

o

:~U)~ .JO:~

-

~

F'I.J: J X~

~°° ° I

~o

I;I ! I-=

"£?

I

I

I

~'~

0

~

X~

--

¢M

{~I

"

0 ~

-o~

IJ.J~

~

0

~

0

~

--

--

N

@J

--

200

K.J. S T R N A T

Bd{t)

I'Op200*C

~"

e,300*C

e"

,/350*

Bd(O )

071I

i

! ,,I , I0

i ill , I02

~ ,~1 , 10 3

EXPOSURET,ME, t

~~ 104

[,OURS]

Fig. 42. Long-term aging behavior of a high-H c 2 - 1 7 magnet at exposure temperatures between 200 and 4 0 0 ° C , permeance p = - 1 . Flux reference is the value after magnetizing and before heating, Ba(t)/Bd(O ) = 1 , loss = 0. (After Ervens 1982a.

Bd(t)

I

1.0

I

ill

J

i

, I I

I

I. I I I

I

I

I _

Bd(O) 0.9 =--~ . .

30/0.C

_

\\~// 0.8

B__. ,.~ - I H

0.7

i

I

ill

I

2-17 \ k ---I-5 \

I III

I0

t

~'all

i

102

i

I

103

EXPOSURE T,ME, t

104

[HOURS]

Fig. 43. C o m p a r i s o n of the aging behavior of sintered S m C o 5 and h i g h - H c 2 - 1 7 magnets at 200 and 300 o C. Permeance p - - 1. (After Ervens 1982a.)

1.0'

I

I

I1[

I

I

I I1

I

I

I I I

I

|

I

Bd(t)

Bd(O)

0.9

~,,~.,,.

H ,~ - 2

2-17 ---

I-5

0.8 0.7

J lie I0

\ I

I Ill I0 2

i

I Ill

I

I I

10 3

04

EXPOSURE TIME, t [HOURS] Fig. 44. C o m p a r i s o n of the aging behavior of sintered S m C o s and high-//c 2 - 1 7 magnets at 400 ° C and two different permeance values, p = - 1 and - 2. (After Ervens 1982a.)

R - C o PERMANENT MAGNETS

~

%

201

<~

E 0

C

',G

0

c,,)

~. ~ E- ~ ~ =

~

202

K.J. STRNAT

rapidly and severely. The comparison at 400 ° C, fig. 44, shows that the 2-17 can still provide a reasonably stable flux even at this high temperature for high permeance values, here for p ~ - 2 , while SmCo 5 magnets are clearly beyond their useful temperature limit for any operating point. It can be concluded that the high-Hc 2-17 magnets permit the highest continuous operating temperatures, certainly to about 300°C, with short heating to 350 ° permitted. The second best REPM for this purpose is well-sintered SmCo 5 with long-term operation at 200 ° C being safe for many purposes and brief periods at about 250 o C permissible.

4.5. Summary of physical properties In addition to the various magnetic properties discussed, which are mostly derived from the B,H- and intrinsic hysteresis loops, an engineer designing a magnetic device or machine usually needs some information on the mechanical, electrical and thermal properties of the magnets. We have paid some attention to these in a qualitative way while weighing the relative merits of bonded versus sintered magnets and of different alloy compositions. We shall now summarize published information regarding those physical quantities frequently required. Table 5 is a composite of data from earlier compilations (e.g., Mildrum and Iden 1975, Nagel 1980, Mildrum et al. 1981) and many commercial product brochures, critically reviewed by the author. There are still obvious gaps in the information and in many cases only single measurements have been reported. On the other hand, there are sometimes many numbers published, especially for the mechanical properties, which cover a very wide range. This may indicate that a listed material category, such as "polymer matrix magnets", indeed covers a broad spectrum of products that have widely varying properties. On the other hand, regarding the mechanical strength of sintered REPM, their behavior is like that of all brittle materials: very much dependent on unavoidable minor flaws and stress risers in the individual sample, and difficult to measure. It is quite possible to find occasionally very low strength values, outside even the wide range specified in the table. Sintered magnets simply should not be made stress-bearing components of a machine. There is still much careful work to be done, both in characterizing and in learning how to control the various nonmagnetic physical properties of interest.

5. Applications of the rare-earth permanent magnets The rare-earth magnets offer much greater energy products than older magnets and also extremely high intrinsic coercive forces, i.e., resistance to demagnetization. The high energy product, especially in dynamic applications, makes possible a dramatic miniaturization of many devices in which the magnet previously was a major part of the volume and weight (e.g., magnetron tubes, magnetic couplings). Their very high coercivity makes the REPM ideally suited for other devices, such as magnetic bearings, which were barely or not at all feasible before the advent of the REPM. And they are also ideal for many types of electric motors, generators and actuators

R-Co PERMANENT MAGNETS

203

(mechanically and electronically commutated DC, stepper, homopolar and synchronous motors; alternators, etc.; rotating and linear), thus accelerating the trend, started by the availability of ferrites and power semiconductors, of returning from the purely electromagnetic (Siemens) to the original permanent magnetic (Faraday) concept of electric machine design. The high cost and relative scarcity of the rare earths, especially Sm, and also of cobalt, long discouraged widespread industrial use. However, imaginative redesign of magnetic circuits to take best advantage of the unique properties of the REPM, and careful analysis of possible cost savings in manufacturing and of the performance of entire systems (including energy savings) convinced designers that the REPM were the best choice in many cases after all. Consequently, their use grew rapidly in the last decade. This caused a fear of a Sm shortage, which has now indeed occurred. This fear fueled the development of the 2-17 magnets and of a variety of Ce and MM-Co-based alloys, and eventually also of the related N d - F e - B magnets which need neither Sm nor Co and have even higher energy products at room temperature (but they have some severe limitations as well). As a consequence, the application engineers now have a wide choice of commercially available R E P M from which to choose the best magnet for a given design, and an even larger variety of types is possible if economics justifies their development. The Nd-Fe-based magnets will obviously displace S m - C o in many present applications where limited temperature capabilities and poor chemical stability are of no concern, and their availability should considerably extend the application range of the REPM because of the broader raw material base. This should relieve the pressure on the Sm supply and price, and make the R E - C o magnets more available for those critical uses where they are the only possible choice. The first conventional devices to apply and benefit greatly from SmCo s were microwave and mm-wave power devices, such as traveling-wave tubes, klystrons, magnetrons and similar tubes in which steady magnetic fields are needed for focusing and guiding an electron beam. These tubes can also take advantage of the still better energy density and high-temperature capabilities of the 2-17 Sin-Co material. The magnets internally temperature-compensated with heavy rare earths, first 1-5 and then 2-17, were also developed for such tubes, but they are finding additional uses in microwave filters, accelerometers and gyros for inertial guidance systems. These are all applications in which S m - C o will have no competition from N d - F e - B or even M M - C o magnets (Rgthwarf et al. 1978). Several unconventional devices have recently been developed in which the REPM must also focus or guide charged particle beams: magnetic lenses, deflection dipoles, multipolar undulators and wigglers for generating synchrotron radiation (Halbach 1985); these, too, seem generally tied to the use of Sm-Co. A large family of highly original permanent magnet field sources, some using a novel "cladding principle" for stray-flux suppression, is now evolving from earlier tube and wiggler structures (Leupold et al. 1987). These constructs are possible only with REPM, Co- as well as Fe-based. (Although for the moment, some of them are solutions looking for a problem.) Also in the category of devices that could only be built with REPM are certain magnetic bearings used in ultra-centrifuges and in turbomolecular vacuum pumps. These and

204

K.J. STRNAT

other magneto-mechanical devices are based on the unique ability of these magnets to produce very strong repulsive forces without getting demagnetized. An exciting application in the line of mechanical force devices is the use of Sm-Co magnets as part of the levitation as well as propulsion system in "magnetic cushion railroads" (Weh 1981). If these should come into commercial use, the quantity demand for REPM, either the Sin-Co used in the prototypes, N d - F e - B , or perhaps MM-Co, will be staggering. A similarly large potential quantity use could develop in magnetic resonance imaging devices, where permanent magnets are competing with superconductors, and the REPM with hard-magnetic ferrites, The REPM are also used extensively in electrobalances, the moving-coil units at the heart of all modern electronic weighing systems, with applications from laboratory to supermarket. It was somewhat surprising that the REPM could also find extensive applications in the field of consumer electronics: for loudspeakers, earphones, microphones, etc., first SmC% and then 2-17. However, this is undoubtedly a market which N d - F e - B will completely take over. In medicine, tiny and even implantable hearing aids have been designed, and artificial heart-assist devices utilize the REPM (Kovacs et al. 1987). The largest market for the REPM in commercial terms, probably even now and certainly in the future, is in small motors and linear or rotary actuators. Computer peripheral devices are now an important part of this: disc-drive motors, recording head positioners, line and dot-matrix printers. Miniature and subminiature REPM motors are also going into the tape or head drives of video cameras and players, into automated optical film cameras, miniature audio tape recorders, electronic typewriters and other business machines, with future applications seemingly unlimited. These small motors are the main application for the new injection-molded 2-17 and N d - F e - B magnets with polymer matrix. Larger REPM motors built to produce very high torques are used in machine tools, industrial robots and in aircraft for moving the flight control surfaces. Very large electric motors of several 100 kVA rating are also under development for a variety of uses, including ship and submarine propulsion. In some of these high-power applications, it is likely that RE-Co will be the magnets of choice because of high operating temperatures and long operating life requirements. References Asti, G., and A. Deriu, 1982, in: Fidler (1982) p. 525. Bachmann, K., 1977, J. Magn. & Magn. Mater. 4,8. Barbara, B., 1978, in: Strnat (1978b) p. 137. Becker, J.J., 1968, J. Appl. Phys. 39, 1270. Becket, J.J., 1969, IEEE Trans. Magn. MAG-5, 211. Becker, J.J., 1976, IEEE Trans. Magn. MAG-12, 965. Benz, M.G., and D.L. Martin, 1970, Appl. Phys. Lett. 17, 176.

Benz, M.G., and D.L. Martin, 1971, J. Appl. Phys. 42, 2786. Benz, M.G., R.P. Laforce and D.L. Martin, 1974, in: AIP Conf. Proc. No. 18/2, eds C.D. Graham and J.J. Rhyne (American Institute of Physics, New York) p. 1173. Bohlmann, M.A., 1976, in: Strnat (1976) p. 31. Bohlmann, M.A., 1987, in: Herget and Poerschke (1987) p. 233. Brandis, H., W. Heimann and H.H. Weigand, eds, 1980, Thyssen Edelstahl Technische Berichte, 1980, Vol. 6, no. 1, Sonderheft

R-Co PERMANENT MAGNETS Dauermagnetwerkstoffe (Thyssen Edelstahlwerke AG, 4150 Krefeld, FRG) ISSN 0340-5125. Burzo, E., 1986, Magneti Permanenti, Vol. 1 (Editura Academiei Republicii Socialiste Romania, Bucharest). Buschow, K.H.J., 1972, J. Less-Common Met. 29, 283. Buschow, K.H.J., 1974, J. Less-Common Met. 37, 91. Buschow, K.H.J., 1980, in: Handbook on Ferromagnetic Materials, Vol. 1, ed. E.P. Wohlfarth (North-Holland, Amsterdam) pp. 297-414. Buschow, K.H.J., 1986, Mater. Sci. Rep. 1, 1-64. Buschow, K.H.J., 1988, in: Handbook on Ferromagnetic Materials, Vol. 4, eds E.P. Wohlfarth and K.HJ. Buschow (North-Holland, Amsterdam) ch. 1. Buschow, K.H.J., and A.M. van Diepen, 1974, Solid State Commun. 15, 903. Buschow, K.H.J., and W.A.JJ. Velge, 1967, J. Less-Common Met. 14, 11. Buschow, K.H.J., and W.A.J.J. Velge, 1969, Z. Angew. Phys. 26, 157. Buschow, K.H.J., W. Luiten, P.A. Naastepad and F.F. Westendorp, 1968, Philips Tech. Rev. 29, 336. Cadieu, F.J., T.D. Cheung and L. Wickramasekara, 1985, J. Appl. Phys. 57, 4161. Cech, R.E., 1974, J. Met. 26, 32. Chino, T., M. Honshima, K. Ohashi and Y. Tawara, 1979, in: Kaneko and Kurino (1979) p. 387. Cremer, R.J., G.W. Reppel and F. Demmel, 1982, in: Fidler (1982) p. 385. Das, D.K., 1969, IEEE Trans. Magn. MAG-5, 214. Das, D.K., and K. Kumar, 1980, IEEE Trans. Magn. MAG-16, 1000. den Broeder, F.J.A., and K.H.J. Buschow, 1972, J. Less-Common Met. 29, 67. den Broeder, F.J.A., and H. Zijlstra, 1976, J. Appl. Phys. 47, 2688. den Broeder, F.J.A., G.D. Westerhout and K.H.J. Buschow, 1974, Z. Metallkd. 65, 501. Deportes, J., D. Givord, R. Lemaire, H. Nagai and Y.T. Yang, 1976, J. Less-Common Met. 41, 273. Domazer, H.-G., 1974, in: Proc. 3rd Eur. Conf. on Hard Magnetic Materials, ed. H. Zijlstra (Bond voor Materialenkennis, Den Haag) p. 140. Domazer, H.-G., and K.J. Strnat, 1976, in: Strnat

205

(1976) p. 348. Doser, M., and L.B. Beach, 1974, IEEE Trans. Magn. MAG-10, 735. Durst, K.-D., and H. Kronmtiller, 1985, in: Strnat (1985) p. 725. Ermolenko, A.S., 1973, in: Proc. Int. Conf. on Magnetism, ICM-73, Vol. I, Part 1 (Nauka, Moscow) p. 231. Ermolenko, A.S., 1976, IEEE Trans. Magn. MAG-12, 992. Ermolenko, A.S., 1982, in: Fidler (1982) p. 771. Ermolenko, A.S., A.V. Korolev and Ya.S. Shut, 1973, Pis'ma v Zh. Eksp. & Teor. Fiz. 17, 449. Ervens, W., 1979, in: Kornfeld (1979) p. 3. Ervens, W., 1982a, in: Tech. Mitteilungen Krupp, Forschungsberichte, Vol. 40, No. 3 (Krupp Gemeinschaftsbetriebe, Fachbticherei, D4300 Essen, FRG) pp. 99-107, ISSN 04949382. Ervens, W., 1982b, in: Fidler (1982) p. 319. Evans, J., P.S. Nyholm, A. Kianvash and I.R. Harris, 1982, in: Fidler (1982) p. 373. Fidler, J., 1981, Philos. Mag. B 46, 565. Fidler, J., ed., 1982, Proc. 6th Int. Workshop on Rare Earth Permanent Magnets and their Applications, Baden (Technical University of Vienna, Austria). Fidler, J., and P. Skalicky, 1982, in: Fidler (1982) p. 585. Fidler, J., P. Skalicky and F. Rothwarf, 1983, IEEE Trans. Magn. MAG-19, 2725. Ghandehari, M.H., and J. Fidler, 1985, IEEE Trans. Magn. MAG-21, 1973. Ghandehari, M.H., R.E. Golden and K.L. McNutt, 1984, IEEE Trans. Magn. MAG-20, 1611. Glardon, R., and W. Kurz, 1978, in: Strnat (1978) p. 504. Glardon, R., and W. Kurz, 1979, Z. Metallkd. 70, 386. Greedan, J.E., and V.U.S. Rao, 1973, J. Solid State Chem. 6, 387. Greinacher, E., and K. Reinhardt, 1982, SeRene Erden, in: Chem. Technologie, 4th Ed., Vol. 2, eds Winnacker and Kiichler (Hanser, Munich) p. 683. Groessinger, R., and J. Liedl, 1981, IEEE Trans. Magn. MAG-17, 3005. Hadjipanayis, G.C., 1982, in: Fidler (1982) p. 609. Hadjipanayis, G.C., 1986, in: Salsgiver et al. (1986) p. 89. Halbach, K., 1985, in: Strnat (1985) p. 123.

206

K.J. STRNAT

Hamano, M., 1987, in: Herget and Poerschke (1987) p. 683. Hamano, M., and S. Yajima, 1978, in: Strnat (1978b) p. 102. Hanna, S.C., and M.S. Walmer, 1985, in: Strnat (1985) p. 309. Harris, I.R., J. Evans and P.S. Nyholm, 1979, British Patent 1 554 384. Herget, C., 1975, in: Kornfeld (1975) p. 3. Herget, C., 1979, in: Kornfeld (1979), Appendix: Phase Diagrams, p. 72. Herget, C., 1987, Metal Powder Rep. 41, 438. Herget, C., and R. Poerschke, eds, 1987, Proc. 9th Int. Workshop on Rare Earth Magnets and their Applications, Bad Soden (Deutsche Physikalische Gesellschaft, Bad Honnef, FRG). Herget, C., H. Kronmtiller and R. Poerschke, eds, 1987, in: Proc. 5th Int. Symp. on Magnetic Anisotropy and Coercivity in Rare Earth-Transition Metal Alloys, Bad Soden (Deutsche Physikalische Gesellschaft, Bad Honnef, FRG). Higuchi, A., and N. Ishigaki, 1979, in: Kaneko and Kurino (1979) p. 355. Hofer, F., 1970, IEEE Trans. Magn. MAG-6, 221. Hoffer, G.I., and K.J. Strnat, 1966, IEEE Trans. Magn. MAG-2, 487. Hoffer, G.I., and K.J. Strnat, 1967, J. Appl. Phys. 38, 1377. Hubbard, W.M., E. Adams and J. Gilfrich, 1960, J. Appl. Phys. 31, 368S. Inomata, K., T. Oshima, T. Ido and M. Yamada, 1977, Appl. Phys. Lett. 30, 669. Inomata, K., T. Oshima and H. Yamamiya, 1978, in: Strnat (1978a) p. 420. Inomata, K., T. Oshima and T. Sawa, 1979, in: Kaneko and Kurino (1979) p. 397. Ishigaki, N., and A. Higuchi, 1978, in: Strnat (1978b) p. 54. James, W., R. Lemaire and F. Bertaut, 1962, C.R. Hebd. Srances Acad. Sci. 225, 896. Jones, F.G., 1987, in: Herget and Poerschke (1987) p. 737. Jones, F.G., and M. Tokunaga, 1976, IEEE Trans. Magn. MAG-12, 968. Kamino, K., and T. Yamane, 1976, in: Strnat (1976) p. 377. Kamino, K., Y. Kimura, T. Suzuki and Y. Itayama, 1973, Trans. Jpn. Inst. Met. 14, 135. Kaneko, H., and T. Kurino, eds, 1979, Proc. 4th Int. Workshop on Rare Earth-Cobalt Permanent Magnets and their Applications, Hakone

(Society of Non-Traditional Technology, Tokyo 105, Japan). Kaneko, H., M. Homma, M. Fukuda and T. Miyazawa, 1977, Digests of 1977 Autumn Meeting of the Japan Institute of Metals (Japan Institute of Metals, Tokyo) p. 225. Katayama, T., and T. Shibata, 1973, Jpn. J. Appl. Phys. 12, 319. Khan, Y., and D. Feldmann, 1973, J. Less-Common Met. 33, 305. Khan, Y., and A.H. Qureshi, 1973, IEEE Trans. Magn. MAG-9, 156. Kianvash, A., and I.R. Harris, 1985, J. Mater. Sci. 20, 682. Kirchmayr, H.R., 1966, IEEE Trans. Magn. MAG-2, 493. Kirchmayr, H.R., and C.A. Poldy, 1979, in: Handbook on the Physics and Chemistry of Rare Earths, Vol. 2, eds K.A. Gschneidner Jr and L. Eyring (North-Holland, Amsterdam) p. 55. Kornfeld, F., ed., 1975, Goldschmidt Informiert 4/75, No. 35 (Th. Goldschmidt AG, Essen, FRG). Kornfeld, F., ed., 1979, Goldschrnidt Informiert 2/79, No. 48 (Th. Goldschmidt AG, Essen, FRG). Kovacs, S.G., M.S. Sweeney, L.E. Ondrovic and R.G. Connar, 1987, in: Herget and Poerschke (1987) p. 181. Kronmtiller, H., 1982, in: Fidler (1982) p. 555. Kumar, K., D. Das and E. Wettstein, 1978, IEEE Trans. Magn. MAG-14, 788. Leamy, H.L., and M.L. Green, 1973, IEEE Trans. Magn. MAG-9, 205. Lee, R.W., E.G. Brewer and N.A. Schaffel, 1985, IEEE Trans. Magn. MAG-21, 1958. Legvold, S., 1980, in: Handbook on Ferromagnetic Materials, Vol. 1, ed. E.P. Wohlfarth (North-Holland, Amsterdam) pp. 183-295. Lemaire, R., 1966, Cobalt 32, 133-140; and 33, 201-211. Leupold, H.A., E. Potenziani II and J.P. Clarke, 1987, in: Herget and Poerschke (1987) p. 109. Li, D., and K.J. Strnat, 1984, J. Appl. Phys. 55, 2103. Li, D., E. Xu, J. Liu and Y. Du, 1980, IEEE Trans. Magn. MAG-16, 988. Li, D., H.F. Mildrum and K.J. Strnat, 1988, J. Appl. Phys. 63, 3984. Liu, S., K.J. Strnat and H.F. Mildrum, 1982, in: Fidler (1982) p. 631. Livingston, J.D., 1973, in: AIP Conf. Proc. No. 10, Part I, eds C.D. Graham and J.J. Rhyne

R-Co PERMANENT MAGNETS (American Institute of Physics, New York) p. 643-657. Livingston, J.D., 1975, J. Appl. Phys. 46, 5259. Livingston, J.D., 1981, in: Prog. Mater. Sci., Chalmers Anniversary Volume, eds J.W. Christian, P. Haasen and T.B. Massalski (Pergamon Press, Oxford) p. 243. Livingston, J.D., 1986, in: Salsgiver et al. (1986) p. 71. Livingston, J.D., and D.L. Martin, 1977, J. Appl. Phys. 48, 1350. Livingston, J.D., and D.L. Martin, 1983, IEEE Trans. Magn. MAG-20, 140. Luborsky, F.E., 1966, J. Appl. Phys. 37, 1091. Martin, D.L., and M.G. Benz, 1972, IEEE Trans. Magn. MAG-8, 35. Martin, D.L., and J.G. Smeggil, 1974, IEEE Trans. Magn. MAG-10, 704. Martin, D.L., J.T. Geertsen, R.P. Laforce and A.C. Rockwood, 1974, in: Proc. l l t h Rare Earth Res. Conf., p. 342. Martis, R.J.J., N. Gupta, S.G. Sankar and V.U.S. Rao, 1978, J. Appl. Phys. 49, 2070. McCaig, M., 1977, Permanent Magnets in Theory and Practice (Wiley, Halsted Press, New York). Melton, K.N., and R.S. Perkins, 1976, J. Appl. Phys. 47, 2671. Menth, A., and H. Nagel, 1978, in: Strnat (1978b) p. 56. Mildrum, H.F., and D.J. Iden, 1975, in: Kornfeld (1975) p. 54. Mildrum, H.F., and K.M.D. Wong, 1976, in: Strnat (1976) p. 35. Mildrum, H.F., A.E. Ray and K.J. Strnat, 1970, in: Proc. 8th Rare Earth Research Conf., Vol. 1, eds T.A. Henrie and R.E. Lindstrom (US Bureau of Mines, Reno, Nevada) p. 21. Mildrum, H.F., M.F. Hartings, K.M.D. Wong and K.J. Strnat, 1974, IEEE Trans. Magn. MAG-IO, 723. Mildrum, H.F., G.A. Graves and Z.A. Abdelnour, 1981, in: Strnat (1981) p. 313. Mildrum, H.F., J.B. Krupar and A.E. Ray, 1983, J. Less-Common Met. 93, 261. Mishra, R.K., G. Thomas, T. Yoneyama, A. Fukuno and T. Ojima, 1981, J. Appl. Phys. 52, 2517. Mitchell, I.V., ed., 1985, Proc. Workshop on N d - F e Permanent Magnets, Brussels, 1984 (Commission of the European Communities, Brussels, Belgium). Nagel, H., 1976, in: AIP Conf. Proc. No. 29, eds J.J. Becker et al. (American Institute of

207

Physics, New York) p. 603. Nagel, H., 1979, J. Appl. Phys. 50, 1026. Nagel, H., 1980, in: Brandis (1980) p. 5. Nagel, H., and A. Menth, 1975, in: Kornfeld (1975) p. 42. Nagel, H., H.P. Klein and A. Menth, 1976, J. Appl. Phys. 47, 3312. Narasiinhan, K.S.V.L., 1981, J. Appl. Phys. 52, 2512. Narasimhan, K.S.V.L., 1986, in: Salsgiver et al. (1986) p. 131. Narasimhan, K.S.Vi., W.E. Wallace, R.D. Hutchens and J.E. Greedan, 1974, in: AIP Conf. Proc. No. 18, Part 2 (American Institute of Physics, New York) p. 1212. Narasimhan, K.S.V.L., M.G.H. Wells and D.V. Ratnam, 1978, J. Appl. Phys. 49, 2072. Nassau, K., L.V. Cherry and W.E. Wallace, 1960, J. Phys. & Chem. Solids 16, 131. Nesbitt, E.A., and J.H. Wernick, 1973, Rare Earth Permanent Magnets (Academic Press, New York). Nesbitt, E.A., J.H. Wernick and E. Corenzwit, 1959, J. Appl. Phys. 30, 365. Nesbitt, E.A., H.J. Williams, J.H. Wernick and R.C. Sherwood, 1962, J. Appl. Phys. 33, 1674. Nesbitt, E.A., R.H. Willens, R.C. Sherwood, F. Buehler and J.H. Wernick, 1968, Appl. Phys. Lett. 12, 361. Nezu, T., M. Tokunaga and Z. Igarashi, 1979, in: Kaneko and Kurino (1979) p. 437. Oesterreicher, H., 1984, J. Less-Common Met. 99, L-17. Ojima, T., S. Tomizawa, T. Yoneyama and T. Hori, 1977, Jpn. J. Appl. Phys. 16, 671. Ormerod, J., 1985, J. Less-Common Met. 111, 49. Ostertag, W., and K.J. Strnat, 1966, Acta Crystallogr. 21, 560. Paladino, A.E., N.J. Dionne, P.F. Weihrauch and E.C. Wettstein, 1975, in: Kornfeld (1975) p. 63. Pan, X., W. Ho and C. Yu, eds, 1983, Proc. 7th Int. Workshop on Rare Earth-Cobalt Permanent Magnets and their Applications, Beijing (China Academic Publishers, 137 Chaonei Dajie, Beijing, China). Perkins, R.S., and S. Strgssler, 1977, Phys. Rev. B 15, 477. Perry, AJ., 1977, J. Less-Common Met. gl, 153. Rabenberg, L., R.K. Mishra and G.J. Thomas, 1982a, J. Appl. Phys. 53, 2389. Rabenberg, L., R.K. Mishra and G.J. Thomas, 1982b, in: FiNer (1982) p. 599.

208

K.J. STRNAT

Ray, A.E., 1974, Cobalt (1974.1), p. 13. Ray, A.E., 1984, IEEE Trans. Magn. MAG-20, 1614. Ray, A.E., 1986a, in: Salsgiver et al. (1986) p. 105. Ray, A.E., 1986b, in: Proc. Rare Earth Permanent Magnet Workshop, 27-28 January 1986, ed. J. Hurt (US Army Research Office, Research Triangle Park, NC, USA). Ray, A.E., and K.J. Strnat, 1972, IEEE Trans. Magn. MAG-8, 861. Ray, A.E., K.J. Strnat and D. Feldmann, 1964, in: Proc. 3rd Rare Earth Research. Conf., ed. K. Vorres (Gordon and Breach, New York) p. 443. Ray, A.E., A.T. Biermann, R.S. Harmer and J.E. Davison, 1973, Cobalt (1974- 4), 90. Ray, A.E., W.A. Sofia, J.R. Blachere and B. Zhang, 1987, IEEE Trans. Magn. MAG-23, 2711. Rinaldi, S., and L. Pareti, 1979, J. Appl. Phys. 50, 7719. Rothwarf, F., H.A. Leupold and L.J. Jasper, 1978, in: Strnat (1978) p. 255. Rothwarf, F., Y. Tawara, K. Ohashi, J. Fidler, P. Skalicky, R. GfiSssinger, H. Kirchmayr, S. Liu and K.J. Strnat, 1982, in: Fidler (1982) p. 567. Sagawa, M., S. Fujimura, H. Yamamoto, Y. Matsuura and K. Hiraga, 1984, IEEE Trans. Magn. MAG-20, 1584. Sakamoto, Y., S. Kojima, K. Kojima, T. Ohtani and T. Kubo, 1979, J. Appl. Phys. 50, 2355. Salsgiver, J.S., K.S.V.L. Narasirnhan, P.K. Rastogi, H.R. Sheppard and C.M. Maucione, eds, 1986, in: Syrup. Proc. on Soft and Hard Magnetic Materials and Applications (American Society for Metals, Metals Park, OH, USA) ISBN 0-87170-234-7. Satoh, K., K. Oka, J. Ishii and T. Satoh, 1985, IEEE Trans. Magn. MAG-21, 1979. Savitskii, E.M., and V.F. Terekhova, 1975, Physical Metallurgy of Rare Earth Metals (Academy Nauk USSR, Moscow), in Russian. Schaller, H.G., R.S. Craig and W.E. Wallace, 1972, J. Appl. Phys. 43, 3161. Schneider, G., E.-Th. Henig, H.L. Lukas and G. Petzow, 1985, J. Less-Common Met. 110, 159. Schweizer, J., 1972, in: Technical Report AFMLTR-72-82 (US Air Force Materials Laboratory, Dayton, OH) ch. II. Schweizer, J., K.J. Strnat and J.B.Y. Tsui, 1971, IEEE Trans. Magn. 1VIAG-7, 429.

Senno, H., and Y. Tawara, 1974, IEEE Trans. Magn. MAG-10, 313. Shimoda, T., 1987, in: Herget and Poerschke (1987) p. 651. Shimoda, T., K. Kasai and T. Teraishi, 1979, in: Kaneko and Kurino (1979) p. 335. Shimoda, T., I. Okonogi, K. Kasai and K. Teraishi, 1980, IEEE Trans. Magn. MAG-16, 991. Shimoda, T., E. Natori and C. Tomita, 1985, in: Strnat (1985) p. 297. St~iblein, H., 1989, in: Handbook on Ferromagnetic Materials, Vol. 5, eds E.P. Wohlfarth and K.H.J. Buschow (North-Holland, Amsterdam). Stoner, E.C., and E.P. Wohlfarth, 1948, Philos. Trans. R. Soc. London A 240, 599. Strnat, K.J., 1967, Cobalt 36, 133. Strnat, K.J., 1970, IEEE Trans. Magn. MAG-6, 182. Strnat, K.J., 1972a, in: AIP Conf. Proc. No. 5, Part 2, eds C.D. Graham Jr and J.J. Rhyne (American Institute of Physics, New York) p. 1047. Strnat, K.J., 1972b, IEEE Trans. Magn. MAG-8, 511. Strnat, K.J., ed., 1976, Proc. 2nd Int. Workshop on Rare Earth-Cobalt Permanent Magnets and their Applications, Dayton (University of Dayton, OH, USA). Strnat, K.J., ed., 1978a, Proc. 3rd Int. Workshop on Rare Earth-Cobalt Permanent Magnets and their Applications, San Diego (University of Dayton, OH, USA). Strnat, K.J., ed., 1978b, Proc. 2nd Int. Symp. on Coercivity and Anisotropy in RE-TM alloys, San Diego (University of Dayton, OH, USA). Strnat, K.J., 1978c, J. Magn. & Magn. Mater. 7, 351. Strnat, K.J., 1979, in: Kaneko and Kurino (1979) p. VIII. Strnat, K.J., ed., 1981, Proc. 5th Int. Workshop on Rare Earth-Cobalt Permanent Magnets and their Applications, Roanoke, VA (University of Dayton, OH, USA). Strnat, K.J., 1983, Report No. NASA CR-168178 (DOE/NASA/0189-83/2) (NASA, Lewis Research Center, Cleveland, OH). Strnat, K.J., ed., 1985, Proc. 8th Int. Workshop on Rare-Earth Magnets and their Applications, Dayton (University of Dayton, OH, USA). Strnat, K.J., 1986, in: Salsgiver et al. (1986) p. 63.

R-Co PERMANENT MAGNETS Strnat, K.J., and G.I. Hoffer, 1966, Report AFML-TR-65-446 (US Air Force Materials Laboratory, Dayton, Ohio). Strnat, K.J., and A.E. Ray, 1975a, in: Kornfeld (1975) p. 47. Strnat, K.J., and A.E. Ray, 1975b, AIP Conf. Proc. No. 24 (American Institute of Physics, New York) p. 680. Strnat, K.J., and A. Tauber, 1983, J. Less-Common Met. 93, 269. Strnat, K.J., and J.B.Y. Tsui, 1970, in: Proc. 8th Rare Earth Research Conf., Vol. I, eds T.A. Henrie and R.E. Lindstrom, p. 3. Strnat, K.J., G.I. Hoffer and A.E. Ray, 1966a, IEEE Trans. Magn. MAG-2, 489. Stmat, K.J., G.I. Hoffer, W. Ostertag and J.C. Olsen, 1966b, J. Appl. Phys. 37, 1252. Strnat, K.J., G.I. Hoffer, J.C. Olsen, W. Ostertag and J.J. Becker, 1967, J. Appl, Phys. 38, 1001. Strnat, K.J., K.M.D. Wong and H. Blaettner, 1976a, in: Proc. 12th Rare Earth Research Conf., ed. C.E. Lundin (University of Denver, CO, USA) p. 31. Strnat, K.J., A.J. Kleman and H. Blaettner, 1976b, in: Strnat (1976) p. 387. Strnat, R.M.W., and S. Liu, 1982, in: Fidler (1982) p. 469. Strnat, R.M.W., and H.-L. Luo, 1982, in: Fidler (1982) p. 457. Strnat, R.M.W., S. Liu and K.J. Stmat, 1982, J. Appl. Phys. 53, 2380. Strnat, R.M.W., J.P. Clarke, H.A. Leupold and A. Tauber, 1987a, J. Appl. Phys. 61, 3463. Strnat, R.M.W., M. Doser and J.P. Clarke, 1987b, in: Herget and Poerschke (1987) p. 675. Suzuki, S., I. Okonogi and K. Kasai, 1978, in: Strnat (1978a) p. 438. Suzuki, T., T. Yamane, K. Kamino, Y. Hasegawa, M. Hamano and S. Yajima, 1979, in: Kaneko and Kurino (1979) p. 325. Tanigawa, S., Y. Hara, E.J. Lavernia, T.S. Chin, R.C. Handley and N.J. Grant, 1986, IEEE Trans. Magn. MAG-22, 746. Tatsumoto, E., J. Okamoto, H. Fujii and C. Inoue, 1971, Suppl. J. Phys. (France) 32, C1-550. Tawara, Y., and T. Chino, 1979, in: Kornfeld (1979) p. 10. Tawara, Y., and H. Sermo, 1968, Jpn. J. Appl. Phys. 7, 966. Tawara, Y., and H. Senno, 1973, Jpn. J. Appl. Phys. 12, 761. Tawara, Y., and H. Senno, 1976, in: Strnat (1976) p. 340.

209

Tawara, Y., and K.J. Strnat, 1976, IEEE Trans. Magn. MAG-12, 954. Tawara, Y., K. Ohashi, S. Iwasaki and N. Kobayashi, 1985, in: Strnat (1985) p. 287. Taylor, R.J., and D.P. Wainwright, 1976, in: Stmat (1976) p. 364. Tenzer, R., 1969, Applied Magnetics, Vol. 16, No. 1, Company Journal of Indiana General Company [now IG Technology] (Valparaiso, IN, USA). Velge, W.A.J.J., and K.H.J. Buschow, 1968, J. Appl. Phys. 39, 1717. Velicescu, M., 1985, J. Less-Common Met. 111, 71. Vein, E.M.T., and E.C. Subbarao, 1986, J. Mater. Sci. 21, 3387. Vogel, R., and W. Frilling, 1947, Metallforschung 2, 97. Walkiewicz, J.W., E. Mortice and M.W. Wong, 1983, IEEE Trans. Magn. MAG-19, 2053. Wallace, W.E., 1973, in: Rare Earth Intermetailics (Academic Press, New York) pp. 112-195. Wallace, W.E., H. Fujii, M.V. Satyanarayana and S. Hirosawa, 1982, in: Fidler (1982) p. 537. Web, H., 1981, in: Strnat (1981) p. 113. Weik, H., and K. Strnat, 1965, in: Landolt-Bornstein, 6th Ed., Vol. IV, section 2c (Springer, Berlin) p. 607. Wells, M.G.H., and K.S.V.L. Narasimhan, 1979, in: Kornfeld (1979) p. 15. Westendorp, F.F., 1970, Solid State Commun. 8, 139. Westendorp, F.F., and K.H.J. Buschow, 1969, Solid State Commun. 7, 639. White, R.M., ed., 1985, in: Magnetic Materials Committee Report, National Research Council, National Materials Advisory Board Publication NMAB-426 (National Academy Press, Washington, DC) pp. 11-26. Yajima, S., M. Hamano and H. Umebayashi, 1972, J. Phys. Soc. Jpn. 32, 861. Yoneyama, T., S. Tomizawa, T. Hori and T. Ojima, 1978, in: Strnat (1978a) p. 406. Zarechnyuk, O.S., and P.I. Kripyakevich, 1962, Kristallografiya 7, 543. Zijlstra, H., 1970, J. Appl. Phys. 41, 4881. Zijlstra, H., 1971, J. Phys. (France) 32, C1-1039. Zijlstra, H., 1982, in: Handbook on Ferromagnetic Materials, Vol. 2, ed. E.P. Wohlfarth (North-Holland, Amsterdam) pp. 37-105. Zijlstra, H., and F.F. Westendorp, 1969, Solid State Commun. 7, 857.