Creep-induced magnetic anisotropy of an amorphous Fe80Cr2B14Si4 alloy

Journal of Magnetism and Magnetic Materials 72 (1988) 199-208 North-Holland, Amsterdam

CREEP-INDUCED

MAGNETIC

ANISOTBOPY

199

OF AN AMORPHOUS

Fe,Cr,B,,Si,

ALLOY

L. KRAUS, N. Z~LRUBOV~~,K. ZxkVkTA and P. DUHAJ a Institute of Physics, Czechoslovak Academy of Sciences, 180 40 Prague 8, Czechoslovakia a Institute of Physics, Slovak Academy of Sciences, 842 28 Bratislava, Czechoslovakia Received 9 November 1987

The magnetic anisotropy of an amorphous Fes&rzB,$i, alloy induced by stress-annealing was investigated. Anisotropy of the order lo4 erg/cm3 with the easy axis parallel to the tensile stress can be induced also at temperatures well above the Curie temperature. The anisotropy is fully removable by subsequent annealing without stress. The dependence of the anisotropy constant on various annealing conditions indicates that thermally activated reversible shear transformations of small atomic complexes are responsible for the anisotropy.

1. Introduction

It is well known that large uniaxial magnetic anisotropy can be induced in metallic glasses by annealing under applied stress [l]. In contrast to the field-induced anisotropy, this type of anisotropy can also be produced by annealing above the Curie temperature. It has been shown that the anisotropy induced by stress-annealing is closely related to the viscoelastic strain (creep) [2]. The creep-induced anisotropy has been intensively investigated in amorphous Co, Co-Fe and Co-FeNi based alloys [3-lo]. Both signs of the anisotropy constant K, have been observed, depending on the composition of the alloy and the annealing conditions. By analogy with the creep strain the creep-induced anisotropy has been divided into two contributions: the anelastic K, (recoverable) and the plastic K,, (permanent) ones [2]. Both contributions have been reported to be proportional to the stress (I applied during the stress-annealing. Because the same behaviour has been observed for the anelastic and the plastic strains can and Q,,, respectively, each component of the anisotropy is expected to be proportional to the corresponding component of the creep strain [31

Here C, and Cr,t are constants. For most alloys studied so far K, was negative (i.e. the hard axis was parallel to the direction of the tensile stress applied during annealing) and K,, was positive or zero *. Various exceptions to this rule have been, however, observed [lo-131. It has been pointed out by Jagielinski and Egami [8] that the separation of the two contributions to the creep-induced anisotropy may be difficult. In order to get the proper separation, the full kinetics of the recovery of K, should be studied. The microscopic origin of the creep-induced anisotropy is now being intensively investigated. From the experiments on the anisotropy induced by combined stress- and field-annealing [S,lO] it has been concluded that the field-induced and creep-induced anisotropies are of different origins. While it is generally agreed that the field induced anisotropy is due to the directional chemical short range order, (CSRO) [14], the creep-induced anisotropy has been explained by the topological short range order (TSRO). It has been recently shown by the EDXD method that the mechanical creep induces structural anisotropy in metallic glasses. After stress-annealing an excess of atomic bonds in the direction perpendicular to the ap*

0304-8853/88/$03.50 0 Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

Note: The usual sign convention (K, > 0 for the easy axis parallel to the tensile stress axis) is used in this paper.

200

L. Kraus et al. / Anisotropy of an amorphous Fe,oCr,B,,Si,

plied stress has been observed [9,15]. It is assumed that the bond-orientation anisotropy gives rise to the observed magnetic anisotropy. The magnetic mechanisms of the anisotropy based on the pair ordering and the random anisotropy models were recently discussed by Szymczak [16]. In this paper we report the study of the creepinduced anisotropy of an Fe-rich amorphous alloy. The influence of various annealing parameters on the induction and recovery of the anisotropy is investigated in an amorphous Fe,,Cr,B,,Si, alloy. 2. Experimental techniques Amorphous Fe,&r,B,,Si, ribbon with a typical cross section of 10 X 0.026 mm2 was prepared by the planar flow casting technique [17]. The addition of 2 at% of Cr was used in order to decrease the Curie temperature T, well below the crystallization temperature TX and to increase the resistance to oxidation during annealing in the air. The structure of the as-cast ribbon and of ribbons after some annealing treatments for long times at the highest temperatures used was investigated by X-ray diffraction. The room temperature value 4aM, = 14070 G, the Curie temperature T, 347” C and the crystallization temperature TX= 470°C were determined from the temperature dependence of the magnetization, measured by means of a vibrating sample magnetometer in the magnetic field of 1 kOe at a heating rate of about 1.1 K/mm. The glass transition between 450 and 465°C was observed by a differential scanning calorimeter. The annealing treatments were performed in air at various temperatures up to 435’C. They were divided into three successive steps: (1) pre-annealing (without stress) for a period of t, at the temperature Tp followed by (2) stress-annealing for t, at T, with the applied stress a; finally some samples were (3) stress-relaxed (i.e. annealed without stress) for a period of t, at the temperature T,. Approximately 7 cm long pieces of the ribbon were used for the pre-annealing and stress-annealing treatments. After the stress-annealing the samples were rapidly cooled by blowing air to nearly room temperature before the stress was removed.

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Domain structures of the as-received ribbon and some heat treated samples were observed by means of a JEOL 733 Superprobe SEM on the shiny side of the ribbon. The detailed arrangement is described in ref. [18]. The induced in-plane anisotropy was measured at room temperature by the method of biased transverse susceptibility [19] on circular discs of diameter of 9 mm spark-cut from the ribbon. The disc was magnetized to saturation by a dc bias field H, in the disc plane. A small in-plane ac field h perpendicular to the dc field H, caused the saturation magnetization M, to oscillate around its equilibrium dc direction. The amplitude m of the oscillations can be expressed through the transverse susceptibility xt = m cos(+ - Q/h, where cp is the angle of the dc field H, and 8 the angle of the equilibrium dc direction of M, with respect to a preferential direction in the disc plane. The transverse susceptibility xt was measured by means of a Hartshorn bridge. The anisotropy constants were calculated from the angular dependences of the effective field

For in-plane uniaxial anisotropy with the anisotropy constant K, the effective field HePf is given by [19]

+2~,

cam 2e/~,

COS~(+ - e),

where the angle 8 can be calculated equation sin( $I -

e)

= K, sin 28/M, Ho.

(3)

from the (4)

The angular dependences of He,, for some stress-annealed samples measured at the frequency of 325 Hz and the bias field Ho = 250 Oe are shown in fig. 1. The angle I#J= 0 corresponds to the ribbon axis and coincides with the direction of the applied tensile stress. The curves measured on the same disc before annealing (not shown here) and after annealing without stress (curve u = 0 in fig. 1) are nearly identical and show sharp minima at + = f 90 O. It has been shown that such minima are typical for the shape anisotropy caused by the oriented surface roughness of the ribbon [20]. In

L. Kraus et al. / Anisotropy of an amorphous Fe&rzBI,Si,

alloy

3. Experimental

results

tpre’ t, =Ih

Tpre=T, = 325 OC I -180°

I

-90°

O0

4 1

9o"

180°

@

Fig. 1. Angular dependence of the effective field H,,f measured by the biased transverse suswptibility method after various stress-annealing treatments. The bias dc field is 250 Oe.

order to separate the undesirable shape anisotropy from the intrinsic anisotropy of the material, it was assumed that the central part of the curve was not affected by the shape anisotropy. The anisotropy constant K, was calculated from the central part of the experimental curve (- 60 o < 9 < + 60 ” ) using eqs. (2) and (3). The small random variations observed on the curves reflect the details of the surface roughness and can be the source of an experimental error. The error in the value of K, determined by the former procedure was estimated to be less than 1.8 x lo3 erg/cm3 for the ribbon under investigation.

201

The angular dependences of I&,( +) of all stress-annealed samples are similar to those of fig. 1 with the maximum of H,, at + = 0, which indicates that the anisotropy constant K, is positive. This has been proved also by the observation of the domain structure of some stress-annealed samples. One example of the typical domain structure is shown in fig. 2. The parallel light and dark stripes correspond to the domains with opposite magnetization along the ribbon axis. We can thus conclude that the easy axis of the creep-induced anisotropy in this alloy is parallel to the tensile stress applied during the stress annealing. The dependence of the anisotropy constant K, on the annealing temperature T, at a constant annealing time was investigated for the applied stress u = 720 MPa. Because of the large heat capacity of the furnace it was difficult to change rapidly the annealing temperature. Therefore the pre-annealing was usually done at the same temperature as the stress-annealing (T, = T,). The dependence of K, on the annealing temperature (for 1, = t, = 1 h) together with the temperature dependence of magnetization are shown in fig. 3. For annealing temperatures below 200°C no creep-induced anisotropy was observed. With increasing annealing temperature K, increases showing a maximum at about 400 o C ,well above the Curie temperature T,. With further increase of

Fig. 2. Domain structure of the stress-annealed Fes&r,BI,Si, (T,=r,=375’C; tp=t,=lh, a=720MPa).

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L. Kraus et al. / Anisotropy

of an amorphous Fe,,,Cr2 Bl,Si4

T,f’C) Fig. 3. Creep-induced anisotropy as a function of the annealing temperature and the temperature dependence of magnetization measured at the applied field of 1 kOe.

the annealing temperature K, sharply decreases. At the highest temperature of 435 o C (the asterisk in fig. 3) the ribbon ruptured after t, = 52 min and was immediately cooled down to room temperature. Though the creep strain (estimated from the reduction of the width of the ribbon) was about 5% the induced anisotropy was very low. The kinetics of creep-induced anisotropy was investigated at constant applied stress u = 720 MPa. The dependence of K, on the time of stress-annealing fa at three different temperatures T, is shown in fig. 4. As can be seen, these dependences are linear in log t, with different slopes for different annealing temperatures. At 375OC saturation of K, occurs after about 2.5 h, indicating that practically all available relaxation processes already took place by that time. The maximum annealing time t, = 24 h used here was insufficient to saturate K, at the two lower temperatures. The creep-induced anisotropy can be removed by subsequent annealing without stress (stress-relaxation). An example of the recovery of K, is shown in fig. 5. Three samples stress-annealed for different periods t, at the temperature of 375OC were relaxed at the same temperature. As can be seen, the dependence of recovery is also linear in log t,; but the rate of the decrease is smaller than

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that observed for the induction (left part of fig. 5). In other words, the recovery of K, is much slower than the induction at the same temperature. E.g., the sample stress-annealed for 1 h had to be relaxed more than 24 h in order to remove all the induced anisotropy. The dependence of K, on log t, for the sample stress-annealed for 16.5 h is shifted to higher relaxation times. This means that the recovery of this sample, stress-annealed for a long time, is about 6 times lower than the recovery of the two samples stress-annealed for the short time. This clearly demonstrates that the kinetic of recovery depends on the thermal history of the sample. After 24 h of stress-relaxation about l/3 of the initial value of K, was still present in this sample. In order to speed up the recovery, the relaxation temperature T, was raised to 425” C (right part of fig. 5). Then K, quite quickly decreased to zero. After the last heat treatment (full circle) some traces of crystalline phase were already present in the sample. The kinetics of recovery at other temperatures (275, 325 and 425 ’ C) are similar to that shown in fig. 5. The K,(log t,) dependences are, however, shifted to lower relaxation times for higher temperatures and vice versa.

h

“E 1.5X P al “S l.O-

I

Y=

0.5 t IO3

IO4

IO5

t, (s-2) Fig. 4. Dependence

of the anisotropy stress-annealing.

constant

on the time of

L Kraw et al. / Anisotropy of an amorphous Fe,Cr2B,,Si,

Tp= T,= 375’C

’

I

h.

1.5 -

‘.

‘\o\

\ \

I I I I I I I I

‘0

‘h ”

\ \ \

‘9,

\ \

. ....

9

mm--\

= t,

tP

\

\I

\

\

\

le

\

103

IO4

1

1

10

lo2

ta Fig. 5. Recovery

\

1

t,

of the creep-induced

anisotropy

The slopes of the decrease plotted against the logarithm of rr are nearly the same for all relaxation temperatures in contrast to the kinetics of induction (fig. 4). It has been proved that K, may always be relaxed to zero (within the experimental error), independently of the stress-annealing conditions. In addition we did not observe the change of the sign of K, during the recovery, in contrast to what was found by Nielsen et al. [21]. We can therefore conclude that the creep-induced anisotropy of this alloy is always positive and is fully recoverable. The dependence of K, on the stress a applied during stress-annealing at three different annealing temperatures is shown in fig. 6. As can be seen, for the lowest annealing temperature the dependence is linear within the experimental error up to 1.2 GPa. At higher temperatures a deviation from linearity is observed. The higher the annealing temperature, the lower the stress at which the deviation occurs. The influence of the stress u applied during the stress-annealing on the kinetics of recovery was investigated on two samples annealed for the same time (t, = t, = 1 h) at the same temperature ( Tp =

O

0

.

IO3

\

\ \

1

h\

‘,”

4

,

‘9 I J/ IO5 ”

+-, I I

4

\

a - 720 MPa 1

I I

\

\

:Ih

\

I

T, - 425 OC

I

\

0.5 -

I

’

\

--__

i

J

’

___---, t

“E Y ?I al 1.0 ‘0 c

I I I/-

T, = 375 ‘C

I----+’

203

alloy

,o

h

I

‘,A

\

u

,,

l( I5

IO4

\\

,

I,

IO3

104

tr

(W on subsequent

annealing

without

stress.

T, = 375°C) under different applied stress (463 and 720 MPa). The recovery of the relative anisotropy K,(t,)/K,(O) is shown in fig. 7. The

2*o

T

P

: T

a

t

0.2

0.4

0.6

0.8

1.0

(I (GPa) Fig. 6. Dependence of the anisotropy constant applied during the stress-annealing.

on the stress (I

204

L. Krauc et al. / Anisotropy of an amorphous Fe80Cr2B,,Si,

1.0 ----Y\ . 0.8 -

s f

\ \

tp= ta= 1 h \

\

\

0.6 -

\ “s,

L_ 2 J

0.4 -

\

‘: \

o cr.463

MPa

0 a ~720

MPa_

\o \ O’e o‘\

0.21

\n

\

:\

10

about 90% of its saturation value but K, of the long pre-annealed sample was about 4 times lower. However, after t, = 24 h nearly the same values of K, were obtained for the two differently pre-annealed samples. This result indicates that the preannealing slows down the kinetics of induction but practically does not affect the saturation value of K”.

T,, = T, = 375°C \

IO2

IO3

IO4

alloy

i

4. Discussion

\ \ n, IO5

tr (set) Fig. 7. Recovery of the anisotropy induced by stress-annealing under different stress.

kinetics of recovery are nearly the same and the small difference is well within the experimental error. We have also studied the effect of pre-annealing on K,. The pre-annealing generally reduces the anisotropy induced by subsequent stress-annealing. This effect is small at low temperatures but increases with increasing Tp. E.g. after increasing the pre-annealing time t, from 1 to 24 h the anisotropy K, induced by stress-annealing (t, = 1 h, T, = Tp, u = 720 MPa) decreased by only about 10% for Tp = T, = 275” C but by about 65% for Tp = T, = 375 o C. Similarly, increasing the pre-annealing temperature from 375 to 425 o C (keeping the pre-annealing time t, = 1 h constant) reduced the anisotropy K, induced by stress-annealing at T, = 375OC (t, = 1 h, u = 720 MPa) by about 66%. There may be two possible reasons for the observed decrease of K,: either the decrease of the saturation value of K, or the slowing down of the kinetics of induction. In order to decide between these two explanations, the kinetics of induction after two different pre-annealing treatments were compared. The pre-annealings and stress-annealings were performed at Tp = T, = 400 o C. The first series of samples was pre-annealed for t, = 1 h, the second for t, = 24 h. After 1 h of stress-annealing (u = 463 MPa) the induced anisotropy of the shortly pre-annealed sample reached already

The positive creep-induced anisotropy (easy ribbon axis) observed in some amorphous alloys has been accounted for either by the plastic contribution K,, or by the magnetoelastic coupling (in alloys with positive magnetostriction constant) [7]. Our experimental results, however, show that a recoverable positive anisotropy can be induced even above T,, which excludes both the magnetoelastic origin as well as the plastic contribution to the anisotropy. The plastic contribution to the creep-anisotropy is zero (within the experimental error) in this alloy even when the plastic strain is quite large. 4.1. Anelastic strain The anelastic part of the creep-induced anisotropy is closely related to the anelastic strain. According to the present microscopic models of anelasticity in metallic glasses the anelastic strain occurs due to local shear transformations of some small atomic complexes (“shear sites”) [22,23,15]. The shear transformations take place by thermally activated jumps which are biased in one direction by the applied stress. This produces a delayed macroscopic strain of the material. It is assumed that the individual shear transformation results in perturbation of the local short range order which contributes to the macroscopic structural anisotropy. Using the model of local shear transformations and the transition state theory, the following equation for the anelastic strain E, has been obtained [22,24] (5)

L. Kraus et al. / Anisotropy of an amorphou FeBoCr2 BIdSi,

with the saturation strain = (3) -1’2&~&,

Ll(~)

sinh(r~,W2kT,),

(6)

where N, is the density of shear sites of volume L?,, y,, the local shear strain produced by a single shear jump and 7 = a/o the equivalent shear stress. The product y,&, is called the activation volume. The relaxation function is given by

alloy

205

If the spectrum of activation energies is broad, which is usually true in the case of amorphous solids, then the exponential functions in eqs. (7) and (10) can be replaced to a good approximation by step functions [22,26] at energies E, = kT, ln[2v,t,

cosh( Ty,,Qo/2kT,)]

and E, = kT,, ln(2v,t,),

= j%E)[I

- exp(-v,(E)&)]

0

dE,

(7)

where q(E) is the distribution function of activation energies E and the relaxation frequencies v,(E) obey an Arrhenius type equation v,(E) = 2v, cosh( ry&/2kT,)

(8)

(I2b)

respectively. Then we get for the relaxation functions g and g’ g = iEa4(E) g’ =

exp( - E/kT,),

(124

JE,Eaq(E)

dE, dE

g’=O

03a) for E, < E,, for E,< E,.

(13b)

where v, is the jump attempt frequency (of the order of the Debye frequency - 1012 s-l). The cash factor in eq. (8) appears due to the asymmetry of the double well potential produced by the applied stress. Note that these equations predict: (1) the nonlinear dependence of can(~) on the stress u and (2) the speeding up of the kinetics of Ean with increasing applied stress. The kinetics of recovery of can can be described in a similar way by the relaxation function g’

Further simplification of eqs. (13) can be obtained by assuming a box-type distribution function q(E) between the energies E, < E,. Then we get

em@,+ tr) =e&c)g’(c,

g’=g

for E,< E,,

g’=g-(E,-E,)/(E2-El)

forE,
g’=o

forE,orE,
T,, t,, T,, t,).

(9)

It has to be taken into account that only those shear sites which have been activated during the stress annealing treatment can take part in the stress relaxation [22]. Then the relaxation function g’ can be written in the form g’(c, =

T,, I,, r,, 1,)

/%E){l - exp[-v,(E)&] 0

Xexp[

-v,(E)t,]

dE,

>

for E,d E,,

g=(E,-E,)/(E,-E,)

forE,
g=l

for E,< E,, (144

and

(14b) As can be seen, eqs. (14) give linear log t kinetics for both induction and recovery of z,. 4.2. Creep-induced anisotropy

00)

where the relaxation frequencies v, now obey the equation v,(E) = 2vo exp( -E/kT,),

g=o

(11)

which we get from eq. (8) by putting the applied stress u = 0.

Because the creep-induced anisotropy of the amorphous Fe,,Cr,B,,Si, alloy is recoverable by subsequent annealing without stress, it can be attributed to the reversible structural changes caused by the anelastic strain. If we assume that the linear relationship between K, and E, holds and eqs. (5) to (14) properly describe the be-

206

L Kraus et al. / Anisotropy of an amorphous Fe8&rzBI,Sil

haviour of can, we can use these equations for the interpretation of our experimental data. Then the kinetics of induction and recovery of K, should be described by the relaxation functions g and g’, respectively, and the saturation value K,(co) should be proportional to ~~(00) given by eq. (6). The dependence of K, on the annealing temperature T, (fig. 3) can be now discussed using eq. (5). At lower temperatures 200°C < T,.< 400°C the value of K, is controlled mainly by the kinetics of the process because the stress-annealing time t, = 1 h. is not sufficient to reach the saturation value K,(cc) as can be seen from fig. 4. The sharp decrease of K, at temperatures T, above 400°C can be explained either by the decrease of K,(cc) or by significant slowing down of the kinetics. As the experiments with various pre-annealing conditions indicate, this sharp decrease of K, is caused rather by the pre-annealing at temperatures T,( = T,) near the glass transition temperature Tg than by the stress-annealing itself. This viewpoint is also supported by the remarkable decrease of K, when the pre-annealing temperature Tp was increased from 375 to 425 o C. The kinetics of K, induced by stress-annealing (fig. 4) can be well fitted by the linear log t, dependence predicted for the box-type distribution of activation energies (eq. (14a)). The parameters Y,,, E,, E2 and y&, could in principle be obtained from these dependences. But it can be seen at the first sight that the dependences of fig. 4 cannot be fitted by a unique set of the parameters. The slope of the K,(log t,)/&(m) dependence should be, according to eqs. (12a) and (14a), proportional to T,. The experimentally observed slopes, however, show much stronger dependence on T,. The kinetics of recovery of K,, agree better with the theoretically predicted relaxation function g’ (eq. (14b)). The slope of the decay of K,(log tr) does not depend on the time of stressannealing t, (fig. 5) and the temperature T, (not shown here). The decay of K,( t,)/&(O) does not depend on the applied stress u in accordance with the theory. According to eq. (12b) the slope of K,(log t,) should increase proportionally to T,. In fact, rather a slight decrease is observed. This can be explained by the influence of irreversible

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structural relaxation on the kinetics of recovery. The shift of the K,(log tr) dependence to higher relaxation times for the sample stress-annealed for long time (upper curve in fig. 5) can be explained in the same way. The irreversible relaxation interferes with determining the parameters v,, E, and E2 from the experimental data (especially the value of Y,,). If we assume v,, = 1012 s-l then we obtain for the two lower curves in fig. 5 the minimum and maximum activation energies of approximately E, = 1.7 eV and E2 = 2.3 eV, respectively. These values are in a good agreement with the activation energies of other reversible phenomena in amorphous alloys but are a bit higher than those found for the creep-induced anisotropy in Co-rich alloys [7,8]. In the long-time stress annealed sample the spectrum of activation energies is shifted by about 0.1 eV to higher values. The linear dependence of the creep-induced anisotropy on the applied stress u observed at the annealing temperature T, = 275 o C (fig. 6) is in accordance with the results on Co-rich amorphous alloys [2,3,6-81. However, the nonlinear behaviour at higher temperatures has not been observed so far. As can be seen from eqs. (6) and (8), there can be two reasons for such behaviour. The first one is the sinh-dependence of the saturation value of E,(M). The second one is the increase of the relaxation frequencies v, with the increasing applied stress. At intermediate annealing times t,, where the saturation value of K, is not reached (which is the case of fig. 6), both these causes can contribute to the nonlinear dependence K,(a). However, the stress dependence of K, at T, = 375O C can be fitted by a sir&dependence quite well (the dotted curve in fig. 6). It is probably because with the stress annealing time t, = 1 h, about 75% of K,(co) is already reached at this temperature (see fig. 4). From this fitting we can estimate the activation volume y,,Q, = 87 A3 for T, = 375 o C. This value is in good agreement with the value of 120 A3 found for VITROVAC 6025 by Jagielinski and Egami [8] from the stress dependence of the spectrum of activation energies. Because with increasing annealing temperature the deviation from the linear dependence occurs at lower stress levels, we conclude that the activation volume is not constant but increases quite steeply

I. Kraus et al. / Anisotropy of an amorphous Fe&r,B,,Si,

T,. This is in agreement with the results obtained by Taub [27] from the homogeneous plastic flow measurements near Tg. The relaxation, basically reversible in nature, often shows some irreversibility [25]. In the theory of section 4.1, it was supposed that the anelastic behaviour is purely reversible. However, as the dependence of the kinetics of recovery of Ku on the thermal history of the sample and the influence of the pre-annealing on K, indicate, the irreversible structure relaxations have to be taken into account. The irreversible structural changes at temperatures below Tg can manifest themselves basically in two different ways [25]: (1) as a decrease of the density of relaxation centers N, and (2) as a change of the distribution of their activation energies q(E). Concerning the origin of the shear sites responsible for the anelastic relaxation two different explanations were proposed. According to the model by Argon and Kuo [22,24], based on the free volume theory, the shear sites are closely related to the free volume elements where the atoms have reduced coordination, which permits them to undergo the local shear transformations. Contrary to this model Egami and Vitek [23] suggest that the local shear transformations are connected with the reorientations of local shear stress fluctuations (T-defects) which are essentially independent of the free volume sites. While the excess free volume is reduced by the structural relaxation below Tg, the computer simulation shows that the spectrum of T-defects remains practically unchanged up to the temperatures near the melting point [23]. In principle it can be decided between these two models by the different responses of the anelastic behaviour to pre-annealing treatments. In fact, the free volume model predicts a considerable reduction of the number of shear sites N, on pre-annealing, which means a decrease of K,,(m). On the other hand the T-defect model predicts the anelasticity nearly independent of structural relaxation [28]. Because the irreversible structural changes are relatively rapid around Tg, pre-annealings at these high temperatures are more convenient for this purpose. Our experimental results show that the creepinduced anisotropy can be considerably reduced with

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201

by pre-annealing near Tg_ The decrease is, however, caused primarily by slowing down of the kinetics of induction, i.e. by the increase of the activation energies E. The density of shear sites No remains practically unchanged. The same result has been reported by Haimovich et al. [9] for VITROVAC 6025. This conclusion seems to rule out the free volume origin of the shear sites. However, the observed increase of the activation energies for the local shear transformations does not support the simple model based on T-defects either.

5. Conclusion The biased transverse susceptibility method was used to study the creep induced anisotropy of the amorphous Fe,,Cr,B,,Si, alloy. The influence of various pre-anneal, stress-anneal and stress-relaxation heat treatments on the anisotropy constant K, was investigated. Our experimental results can be summarized as follows: (1) The easy axis of the induced uniaxial anisotropy is always parallel to the tensile stress applied during stress-annealing. (2) The anisotropy constant approaches its saturation value if the stress-annealing lasts for a sufficiently long time. (3) The anisotropy is fully removable by subsequent annealing without stress. (4) The kinetics of induction and recovery indicate a broad distribution of activation energies. (5) The dependence of K, on the applied stress is nonlinear at high annealing temperatures. (6) The irreversible structural relaxation changes the spectrum of activation energies but does not influence essentially the density of relaxation centres. The experimental results can be well understood in terms of the model of anelasticity by Argon and Kuo [22,24] based on the idea of local shear transformations of small atomic complexes. However, the proper explanation of the microscopic origin of the shear sites requires further theoretical and experimental work. The question about the possible plastic contribution to the creep-induced anisotropy, reported

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L. Kraus et al. / Anisotropy of an amorphous Fe,,Cr, B, *Si4 alloy

by some authors, is still open. The results by Haimovich et al. [9] as well as our results show that the kinetics of induction and recovery at the same temperature may be quite different. Therefore the time t, necessary for complete recovery of the anelastic part of the anisotropy can be considerably longer than the time of stress-annealing t,. In order to decide whether the plastic strain contributes to the anisotropy, the full kinetics of the recovery should be studied.

Acknowledgements

The authors would like to thank Dr. V. Kamberskjr for his helpful discussions and critical reading of the manuscript.

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