Fe multilayers

~

journalof magnetism

, 4 4 4 and

magnetic

ELSEVIER

Journal of Magnetism and Magnetic Materials 136 (1994)251-259

,~

materials

Magnetic and magneto-optical properties of Tb/Fe multilayers J. P o m m i e r

a, j . p . J a m e t

a, j . F e r r 6 a,., p. H o u d y

b, p . B o h e r

b, F. P i e r r e b

a Laboratoire de Physique des Solides, associ~ au CNRS, B~t. 510, Universit£ Paris-Sud, 91405 Orsay, France b Laboratoire d'Electronique Philips (LEP), 22 avenue Descartes, 94453 Limeil Brevannes C~dex, France

Received 13 January 1994; in revised form 8 March 1994

Abstract We report on magnetic and magneto-optical properties of [Tb/Fe], multilayers. The dependence on Tb and Fe thicknesses and on the stacking parameter n have been studied. For a thick sample (n = 40), we confirm that the ferrimagnetic or compensation temperatures and the magneto-optical effects are close to those deduced previously for T b x - F e a _ x metallic alloys having the same nominal composition. Thinner samples (n < 40) exhibit slightly different magnetic properties. In particular, a transition from in-plane to out-of-plane anisotropy for low values of n is evidenced and discussed. Magnetization reversal dynamics are studied and analyzed from magnetic after-effect measurements and microscopic domain imaging. Typical values of Barkhausen volumes are deduced.

I. Introduction Rare earth/transition metal multilayers display strong perpendicular magnetic anisotropy [1-5]. Magnetic and magneto-optical investigations are still required to understand the origin of the anisotropy, exchange interaction, magnetic domain structure and to check the ability of these multilayers to be used as magneto-optical recording media. The present study is devoted to the T b / F e system in which the ultrathin terbium and iron layers are ferrimagnetically coupled. This comes from the relative sign and magnitude of the involved exchange integrals: JFe-Ve > --JTb-Fe > Ja'b-a'b. AS for T b x - F e l _ x metallic alloys, a compensation temperature can exist within

* Corresponding author.

a limited range of the ratio p = eTb/eFe between terbium and iron layer thicknesses. The origin of the perpendicular magnetic anisotropy in multilayers is generally attributed to the presence of a T b - F e intermixing state at the interfaces and to large magnetocrystalline effects for terbium ions [4-6]. The competition between the magnetic exchange interactions often leads to noncollinear or conical spin structures [5,6]. A large magnetic field is then necessary to obtain the fully saturated out-of-plane magnetization [7]. Thermal evaporation of metals in ultrahigh vacuum or their sputtering in a high residual vacuum allow the preparation of good samples exhibiting similar magnetic properties [6-11]. We first recall the correlation between crystallographic and magnetic properties by collecting all our previously reported data for a wide set of T b / F e multilayers prepared by rf sputtering [1013]. The magnetic properties of T b / F e multilay-

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SSDI

J'. Pommier et aL /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

252

ers and Tbx-Fel_ x alloys are compared, and the influence of the number of stacking T b / F e periods on their magnetism is discussed. Finally, the dynamics of the magnetization reversal under field are investigated from magnetic after-effect studies and are correlated with the evolution of the magnetic domain structure directly observed by magneto-optical Faraday microscopy [13]. Imaging data are found to be consistent with previous reports on other T b / F e multilayers [14].

2. Elaboration and sample structure

Our n-period [Tb(exb)/Fe(eve)] ~ multilayers were deposited by reactive diode rf sputtering (normal pressure 10 -8 Torr, deposition pressure 5 × 10 -3 Torr of Ar) at room temperature, with a deposition rate of 0.05 n m / s [10]. These structures were grown on a 10 nm thick Si3N 4 buffer deposited on float glass or Si(111) substrates (Fig. 1) and covered by a 10 nm thick Si3N 4 top layer for protection. The silicium substrates were first chemically etched outside the deposition chamber. The structural characterization and the determination of the layer thicknesses of our samples were performed by X-ray diffraction and grazing X-ray reflectometry techniques, while in situ ellipsometry informed on the stacking process and on the interface quality [10]. The main magnetic properties exhibited by our samples are related to the following crystallographic features [10-13]: (i) Fe remains amorphous up to a layer thick-

.......

Si 3N4 Tb Fe

n

~

~

periods:

i.,.,

.,~,,,,.,., , , . . , , . . , , ~ ~ SiaN4

~

~Glassor silicium

Fig. 1. Stackingstructure of the studiedTb/Fe multilayers.

ness of 1.8 nm. This fact is consistent with other data obtained on if-sputtered [7] or thermally evaporated samples in ultrahigh vacuum [6]; (ii) for eTb < 0.35 nm (1 atomic layer (AL) of Tb) the crystallization of iron occurs through Fe bridges inside the terbium layers; and (iii) the interface roughness and Tb-Fe interdiffusion are estimated to be less than 2 AL, so that alloying could be present at the interfaces. Considering multilayers with the same T b / F e building blocks, their magnetic properties do not vary much with the stacking parameter n value for n > 20; thus we report below only our results for a [Tb/Fe]40 multilayer. The different magnetic behaviours for thick and thin multilayers were pointed out previously by Sato [1].

3. Magnetic and magneto-optical properties of thick T b / F e multilayers (n = 40)

3.1. Magnetic anisotropy phase diagram The evaporation and rf sputtering techniques in high vacuum provide T b / F e multilayers having similar crystallographic and magnetic properties [6-12], allowing a direct comparison of the properties of these samples. It is known that T b / F e multilayers can exhibit a strong perpendicular magnetic anisotropy at least twice as large as that of homogeneous Tbx-Fel_ x alloy films with the same nominal composition [9]. Their magnetic anisotropy phase diagram in the (eXb, eFe) plane has been determined recently [12,13] from magneto-optical and Mrssbauer experiments. Our main conclusions can be summarized as follows: - for ultrathin terbium layers (evo < 0.5 nm) the iron structure is polycrystalline (ii) and since the interface roughness is estimated to be about 2AL (iii), bridges between iron layers are indeed present, forcing the magnetization to lie in the sample plane; - in the same context, for ultrathin iron layers (eve < 0.3 nm) complex magnetic properties are found; - considering the change in crystalline structure (i), the magnetic anisotropy varies drastically from an out-of-plane to an in-plane orientation

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

a

'.2

p =o.25

.2 ~

1

--°1

--.2-6i

p = 0.325

b

".1

253

-4

2

-2

--.2-6i

6i

4 i

-4

-2

4i

-i

6i

H/kOe

H/kOc -9_

2i

U--

p =0.44

d

p = 0.55

"°1 o

0

1

-°1

~ --.1

[

--.2_~ ~

_~

?

4

6

H/kOe

H / kOe _

.7-

0.64

p =

-le

p = 0.69

".Z

.1

f

"°1

0

• -.1 -.2~

~ ~

-2

--.1

~ H/kOe

"2- 6

p=O.75

~ -.l

~

"~

i

2

4

6

i

i

i

-"2--6 " - 4 - 2

2

.2 p = 0.97

? 4 6 H/kOe

p=o.83

0 4

6,

~ j p = 1.03

~

'~o .4-2

6,

H/kOe

10 2-6

~

h

H/kOe

.2

~

-.2

.I

".2-6-4-2

-2 H/kOe

.2 •~

~

"""2-6,

-4,

-2

~

f

2,

4,

6,

H/kOe

Fig. 2. Room-temperature polar Kerr rotation (A = 632.8 nm) hysteresis loops of a set of [Tb(eTb)/Fe(1 nm)]a0 samples deposited onto Si (111) for several values of p = eTb/eFe.

254

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

for an iron thickness close to eFe ~ 1.8 nm [11,12], in agreement with other authors [6,7]; - very square hysteresis loops in perpendicular field are obtained for p values (p = ero/eve) lying between 0.6 and 1.1, at least in the 0.7 < eFe < 1.3 nm range [11-13]. The above behaviour is illustrated in Fig. 2, showing the polar Kerr hysteresis loops of a set of [Tb(eTb)/Fe(1 nm)]40 samples when varying ero from 0.25 to 1.03 nm at a fixed iron thickness, eve = 1 nm. It is well known that polar magnetooptical effects probe a combination of out-ofplane Fe and Tb magnetization components. Note that the homogeneity of the present samples was improved in comparison with those previously studied [10,11].

3.2. Magnetic anisotropy: comparison with alloys

0.4

0.6

0.8

1.0

I

I

I

I

F

I

I

0.15

I 0.2 x

I 0.25

p

400 300 ~200

E-

l00 0 0.1

0.3

Fig. 3. Variation of the compensation temperature of Tb x F e l - x alloys (full line, Ref. [16]) and for some T b / F e multilayers as a function of x and p (from expression (2)). Experimental values for multilayers reported by Sato [1] (<)), Bayreuther et al. [7] ([]), Endl et al. [9] (zx), and Yamauchi et al. [18] ( O ) are compared to our present data (A).

To determine the relation between the mean chemical Tb concentration x and the known value of p, we approximated the amorphous structure by a cubic arrangement in each metallic layer. Within this assumption one gets:

variations of Tcompfor thick multilayers with x or p. In this case one can express Tcomp as a function of p:

x = p / ( p + 2.74).

Tcomp(in K)

(1)

The anisotropy remains perpendicular for multilayers down to p = 0.35 (Fig. 2 and Refs. [11,12]). Thus, one unexpectedly finds that T b / F e multilayers and T b x - F e l _ x alloys exhibit out-of-plane magnetic anisotropy for exactly the same range of nominal compositions (11 < x < 29 at% Tb) [15]. As already mentioned by Bayreuther et al. [7], even for eFe < 1.8 nm the T b / F e multilayers become in-plane magnetized or even nonmagnetic for large terbium thicknesses [11].

3.3. Compensation temperature In T b x - F e l _ x metallic alloys the net magnetization vanishes at the compensation temperature Tcomp, which occurs in the range of concentrations 0.133 < x < 0.272, Tcomp varying correspondingly from 0 to 410 K [16] (Fig. 3). The same behaviour also occurs in T b / F e multilayers [1,7,9,11] for eFe up to 1.8 nm. Assuming the validity of expression (1), we plotted in Fig. 3 the

2.56p - 1.07 2.74 + p 103"

(2)

All previously reported experimental values of Tcomp for multilayers [1,7,9,18] lie a little above the curves obtained for the corresponding alloys [16]. This is also true for the room-temperature data given for our [Tb(eTo)/Fe (1 nm)]40 set of samples (Fig. 4a). All possible errors in the determination of the Tb and Fe layer thicknesses ( A p / p ~ 0.05), in the crude evaluation of the mean distance between Tb or Fe ions within each amorphous layer ( A x / x ~ 0.01), and in the scattering of the Tcomp(X) reported values by various authors [1,16,17], have to be taken into account to explain this difference. This indicates that again Tcomp for T b / F e multilayers is slightly underestimated by the proposed expression (2). As expected, fields of less than 7 kOe are not sufficient to saturate the magnetization near Tcomo (for P = 0.75) (Fig. 2g) because of the strong increases in the perpendicular anisotropy and the coercive field [1], due to the cancellation of the net sample magnetization. In conclusion, we confirm that the

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

values of the ferrimagnetic temperature (Tc ~ 400 K) and the compensation temperature are nearly the same for thick T b / F e multilayers and for thick homogeneous T b x - F e t -x alloy films with the same nominal composition [7].

3.4. Magneto-optics in thick Tb / Fe films The specific Faraday rotation per unit length [0 r] of thick T b / F e multilayers is also close to that found for the Tbx-Fea_ x alloys with the same composition. For example, we find [0 F] = 1.72 × 10 5 and 1.8 × 10 5 d e g / c m , respectively, for [Tb(0.5 nm)/Fe(0.65 nm)]86 and [Tb(1.09 n m ) / Fe(1.2nm)]40 multilayers; these values can be compared with 1.7-1.8 × 10 5 d e g / c m reported for the corresponding alloys [16]. The polar Kerr i

5

I

i

I

4~D

~210 0.2

0.4

0.6

0.8

1

1.2

P = erb/% e

0.3

~

I

,

j

~

'

0.2-

255

rotation magnitude depends slightly on the p value (Fig. 4b), but again its magnitude (0.2 ° at 632.8 nm) is close to that obtained for alloys. Indeed, the Kerr rotation can be significantly enhanced when covering the multilayers with a Si3N 4 adaptive top layer. For example, the polar Kerr rotation, measured at 820 nm, for a [Tb (0.9 n m ) / F e ( 1 . 2 nm)] 6 sample is enhanced from 0.20 up to 1.5 ° using a 695 nm thick Si3N 4 top layer, in agreement with our magneto-optical calculations developed in the frame of the electromagnetic theory.

4. Magnetic and magneto-optical properties of thin T b / F e multilayers (n < 20) Thin multilayers refer to samples where the number of stacking periods n < 20, the total magnetic multilayer thickness being less than 40 nm. Strong changes in the magnetic properties are expected with decreasing n [1]. In order to estimate the relative contributions of the volume (Kv) and the interfaces ( K s) to the total anisotropy K one can vary the T b / F e period thickness A, while preserving the ratio p = evo/eF~. This obviously affects the interface contribution per unit volume and modifies the anisotropy. Different models have been used to fit the variation of K with thickness [2,5,8]. For ideal multilayers, it is also possible to study the anisotropy from a set of samples having small n values (1 < n < 10), since if one uses a relation similar to that used by Honda et al. [8], we get:

"~ 0.1O "~ 0-

-0.s,h "~ -0.2-0.3 0.2

b)

o~ 0.4

l ~ 0.6

/ , 0.8

, 1

1.2

p = erb / eFe Fig. 4. Room-temperature coercivity (a) and polar Kerr rotation, measured at A = 632.8 nm (b), of [ T b ( e ~ ) / F e (1 nm)]40 deposited onto Si (111), as a function of the reduced terbium layer thickness. The samples exhibit out-of-plane magnetic anisotropy in the hatched region.

where K v is the sum of the magnetocrystalline and magnetostatic contributions. As usual, the first term of expression (3) is negative and tends to align spins in the sample plane. The second term, due to the interface anisotropy, is positive for our system and larger than the first term in thick samples (n ~ ~). However, as displayed by expression (3), when reducing n, the surface term can eventually become smaller than AKv, giving in-plane anisotropy at small stacking parameter values. By means of polar and longitudinal Kerr

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

256

Table 1 V a l u e s of the p a r a l l e l (ncll, rll) and p e r p e n d i c u l a r ( H c ±, r ± ) coercivity and s q u a r e n e s s ratio r of the hysteresis loop, respectively, for a set of [Tb(0.95 n m ) / F e ( l . 1 5 nm)] n samples, with different n u m b e r s of stacking periods n. T h e v a l u e s of H c are d e d u c e d from hysteresis loops m e a s u r e d at a field s w e e p i n g rate of 1 k O e / s . V a l u e s of the C u r i e and c o m p e n s a tion t e m p e r a t u r e s are also r e p o r t e d . n

1

3

5

7

Hcl I (Oe) rbl H c± ( k O e ) r± T c (K) Tcomp (K)

2 1 No

2 1 No -

No . . 0.5 1 328 ~ 150

No No . . 1.1 > 5 1 1 345 362 ~ 240 ~ 330

10

20 No 2.8 1 376 ~ Tc

temperature Tc (Table 1) of the [Tb/Fe]~ Ising multilayers increases with n, according to the variation of the magnetic dimensionality from two to three. Simultaneously, Tcomo is significantly reduced when n varies from 10 to 5. For n = 20, Tcomp lies close to Tc, a direct consequence is the change of sign of the magneto-optical effects between the n = 20 and the n = 7 or 5 samples at room temperature.

5. Magnetization reversal dynamics 5.1. Generalities on dynamics

coupled measurements, we evidenced such a transition from out-of-plane to in-plane anisotropy by decreasing n for the set of [Tb(0.95 nm)/Fe(1.15 nm)]~ samples (n = 20, 10, 7, 5, 3, 1) (see Table 1). Starting from expression (3), for n ~ oo we deduce the critical T b / F e periodicity, A* = 12Ks/Kv I, at which K cancels. In thin samples the transition would occur for n * = 1 / 2 ( 1 A / A * ) . From our previous work [11], A* can be tentatively estimated to be about 2.3 nm (at fixed eFe = 1 nm) or 2.7 nm (at fixed ea-b = 1 nm). These values are consistent with the A* = 2.7 nm deduced from Sato's data for the p = 0.794 sample. For our set of samples, with p = 0.826, we can assume that A* lies in the range 2.4-2.7 nm, giving 2 < n * < 4, which is consistent with our experimental result: 3 < n* < 5 (Table 1). Thus this transition from in-plane to out-ofplane anisotropy, which occurs when increasing the stacking number of T b / F e periods, can be understood without assuming imperfect initial multilayer growth conditions [1]. The first iron layer is already magnetic at room temperature, as we evidenced in the bilayer (n = 1) sample (Table 1). The reduction in the coercive field (Table 1) with decreasing n (for n > 5) is directly correlated to the reduction in the perpendicular anisotropy. As expected, the ferrimagnetic Curie

As already reported for T b x - F e l _ x thick alloy films, the dynamics of the field-induced magnetization reversal process is slow [19-22]. Thus, the coercive field value depends on the field sweeping rate and the magnetic after-effect is often used to investigate dynamics. The magnetic relaxation is then strongly dependent on the applied field value. As mentioned previously, the magnetization reversal dynamic processes in [ T b / F e ] n multilayers [14] are similar to those observed in T b x - F e l _ x thick alloy films. For all our studied [Tb(0.95 nm)/Fe(1.15 nm)]n thin samples with n = 5, 7, 10 and 20, the spin reversal is dominated by domain wall motion, i.e. nucleation takes place only at a few magnetic defect places. Near Tcomo, large circular domains are favoured in such samples since the magnetostatic energy term vanishes at the same time as the net magnetization. This behaviour is illustrated in Fig. 5 for the n = 5 sample deposited on a glass substrate. The time-dependent hysteresis loop (Fig. 5a), the magnetic after-effect (Fig. 5b) and the associated evolution of the magnetic domain structure (Fig. 5c) are presented. Magnetic relaxation curves were obtained using a light beam of about 1 mm in diameter. Magnetic domain imaging was performed by means of Faraday rotation microscopy [23]. 5.2. Phenomenological theory and discussion According to the phenomenological theory for ferroelectrics, first developed by Fatuzzo [24],

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259 0.1

I

~, 0.05

a)

O ..~

domains and their wall propagation mechanism. Its time (t) dependence is then expressed by:

I

I

0.9 kOe/sI ~009 ~Oe/s

M/Ms = - 1 + 2 exp{ -ot[k2R2t 2 + Rt

0

.o

+ ( 2 k 2 - 2k + 1 ) ( 1 - R t - e x p ( - R t ) ) ] } ,

,~-0.05

J

-0-1.1"

J

i

i

-1

-0.5

0

(4)

I

I

0.5

1

1.5

H/kOe 0.l -153 Oe O

b)

R

0

-O.l

I

I

I

I

l0

20

30

40

7 s

t/s

50

Q

Q

10 s

13 s

c) 16

257

s

20

25 s

s

[

I

100

~m

Fig. 5. Magnetization reversal of the [Tb(0.95 n m ) / F e (1.15 nm)] s sample deposited onto glass, studied by Faraday rotation. (a) Magnetic hysteresis loops at two field sweeping rates. (b) Magnetic after-effect measured at different negative fields after premagnetizing the sample in a large positive field ( H = 600 Oe). The switching of the field occurs at time t = 0. (c) Domain dynamics in H = - 8 0 Oe. The spin-down domain (in black) expands inside the original spin-up (white) singledomain state. The elapsed time t is indicated.

and extended by Labrune et al. [21] to magnetic materials, the magnetization reversal can be described taking into account both nucleation of

where M s is the saturation magnetization. The parameter k = v / R r i expresses the relative importance of domain wall motion as opposed to the nucleation rate R in the magnetization reversal process, where r i stands for the initial radius of a nucleation centre. The filling coefficient a = NorrrZ/A depends on the total number of nucleation centres N o over the investigated sample area A. Expression (4) shows that M / M s is a function of a, k and Rt. The characteristic time t o taken to reach an M = 0 demagnetized domain state sample can be deduced from the general equation Rt o = f ( a , k). Then: M / M s = F ( a , k, Rt) = G ( a , k, t / t o ) . (5) For the studied samples, as shown for example in Fig. 5(c), magnetization reversal always initiates from a small number of centres with the same energy barrier for nucleation. Thus, in this case, a can be considered to be independent of H, assuming that Barkhausen volumes for nucleation and propagation are identical. It is then possible to show that k does not depend on the applied field value H [21]. So, at a given temperature, M / M s (expression (5)) is only dependent on the reduced time t/to, irrespective of the field value. For each sample, the room-temperature relaxation data (Fig. 5b) can thus be fitted by a unique M / M s universal law (Fig. 6). The domain nucleation does not necessarily occur during field switching (Fig. 5c), a nucleation lag time ~- is often necessary in these samples to allow wall propagation. This universal law fails at long times (Fig. 6) probably due to the restricted area of the viewing field of our microscope and to the presence of remaining hard magnetic centres. The universality of the behaviour observed for different applied fields is consistent with the narrow distribution of Barkhausen volumes in the samples. As for T b x - F e l _ x alloys, we checked di-

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

258

-~.0

-1 I

0

I

t/to

I

5

Fig. 6. Room-temperature magnetization reversal data of the [Tb(0.95 nm)/Fe(1.15 nm)]5 sample collected for different applied fields given in Fig. 5(b). All data points overlap to give a continuous universal curve according to that calculated ( - - - - - - ) by the proposed relaxation law (expression (5)), using k = 55.1, Rt o = 7.23 and a = 5.6× 10 -6. rectly by imaging that the domain wall velocity varies like v = A e x p [ - H / H o] [21,14]. H 0 values for the different samples are reported in Table 2. This variation is consistent with a thermally activated process for domain wall motion, giving: v = v 0 exp

- (Ep - 2 V B M s H ) ~-~ ,

(6)

where VB stands for Barkhausen volume for propagation and Ep is the typical activation energy for wall propagation. Va and the typically associated Barkhausen length h s have been calculated for all samples, either starting from H 0 deduced from domain wall velocity measurements or derived from the t o ( H ) variation (Table 2). Absolute measurements of the magnetization in our samples were not available, so reasonable indirect estimates of M s have been extracted from T b - F e alloy data [1] for samples having nearly the same compositions and compensation Table 2 Values of the magnetization reversal parameters ( H o is defined from v = A e x p [ - H / H o ] ; VB and h n are the Barkhausen volume and length, respectively) deduced for the [Tb(0.95 nm)/Fe(1.15 nm)]s,7,z0 samples n

H0 (Oe)

VB ( × 10 17 cm3)

Aa ('~)

5 7 20

20-25 60-65 80-95

0.8 -1 0.37-0.4 0.20-0.25

~ 170 ~ 90 ~ 40

temperatures. One finds that M s is almost independent of n: for these experiments, the n = 20, 7 and 5 samples yielded M s = 100, 80 and 100 Oe, respectively. The values of h B (Table 2) agree with those (h B = 30-300 A) previously found in other samples [14], but we have no explanation for the variation of the deduced Barkhausen length with the stacking p a r a m e t e r n. The lag time z necessary to initiate propagation after switching the field varies as exp[H/Ho], in agreement with a usual thermally activated process for nucleation [21]. For example, this lag time is 5.3 s for the n = 5 sample, as determined from extrapolation to zero of the time-dependent domain area (Fig. 5c).

6. C o n c l u s i o n s

A general study of the static and dynamic magnetic properties of T b / F e multilayers has been presented. As already discussed, these multilayers exhibit large perpendicular magnetic anisotropy and very square hysteresis loops over a large range of compositions. It is puzzling to note that most of their magnetic properties are identical to those of Tbx-Fe~_ x alloy films with the same composition, in spite of their strongly different exchange interaction mechanisms and the presence of interfaces. This can be understood by assuming that intermixing between Fe and Tb ions is very important even far from the interface, although the sharpness of interfaces demonstrated by ellipsometry measurements contradicts the former hypothesis. We have demonstrated how the in-plane-outof-plane anisotropy transition can be induced by increasing the stacking p a r a m e t e r n for given values of eTb and eve. Magnetization reversal dynamics can be well understood from direct imaging of the time evolution of the magnetic domain structure. However, the variation of the deduced Barkhausen volumes with n is still not understood. Much more fundamental work remains to be done on very well characterized multilayers with small n values, in order to achieve a better description of the anisotropy and magnetism in these ultrathin structures. Their use

J. Pommier et al. /Journal of Magnetism and Magnetic Materials 136 (1994) 251-259

as magneto-optical recording media is underlined.

Acknowledgements This work has been realized thanks to a grant from the French Ministry of Research and Technology. The authors wish to thank V. Grolier and P. Meyer for helpful discussions.

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