study of amorphous and crystallised fe75b25 by auger electron spectroscopy and x-ray photoelectron...

Applied Surface Science 51 (1991) 139-155 139 North-Holland

Study of amorphous and crystallised Fe75B25 by Auger electron spectroscopy and X-ray photoelectron spectroscopy

S.K. S h a r m a 1 and S. H o f m a n n Max-Planck-lnstitut fiir Metallforschung, lnstitut ]'fir Werkstoffwissenschaft, Seestrasse 92, W- 7000 Stuttgart 1, Germany

Received 21 March 1991; accepted for publication 12 May 1991

Auger electron spectroscopy (AES) in conjunction with argon ion sputtering was used to study the sputtering behaviours of an amorphous Fe75B25 alloy and its corresponding crystallized alloy forms, namely the single-phase metastable Fe3B and a two-phase mixture of a-Fe and Fe2B. Bombardment with the energetic argon ions of energy 1 and 3 keV resulted in the establishment of a boron-rich sputter altered layer in all the alloy forms. The magnitude of enrichment of boron was relatively much higher and similar in the amorphous and the metastable Fe3B alloys than in the alloy consisting of a two-phase mixture of a-Fe and Fe2B. The sputter-induced compositional differences in these alloys have been rationalized in terms of their different structural states and are suggestive of a similar local atomic order in the structures of the amorphous and the corresponding crystallized Fe3B. The binding effects appeared to play a significant role in determining the sputtering behaviours of these alloys and the probable sputtering mechanisms would include cascade sputtering and/or surface segregation. The evidence for the occurrence of bombardment-induced surface segregation is provided by a simple model calculation based on the values of the concentration ratio of iron to boron obtained by using the low energy (47 eV) and the high energy (651 eV) Auger peaks of iron.

Native oxide films formed on these structurally different alloy forms were also characterized using AES and X-ray photoelectron spectroscopy (XPS). The oxide films were enriched in boron, the enrichment being maximum near the oxide/alloy interracial region and incorporated Fe 3+, Fe 2+, and B 3+ species. The oxide films on amorphous Fe75B25 and Fe3B were similar in thickness (3.6 and 3.3 nm) while a relatively thicker film (9.7 mm) was formed on the alloy consisting of a-Fe and Fe2B.

1. Introduction

Bombardment of the surfaces of solids by energetic ions results in compositional changes which can be easily monitored by the application of surface-sensitive techniques such as Auger electron spectroscopy (AES), X-ray photoelectron spectroscopy (XPS) etc. Such bombardment-induced compositional changes arise due to preferential sputtering which is caused by the differences in the sputter yields of the constituent elements in a multicomponent target [1]. The effect of preferential sputtering has been examined in many alloy systems and has also been reviewed in the literature [2-7]. The majority of these investigations have been carried out on crystalline al-

I On leave from Physical Metallurgy Division, Bhabha Atomic Research Centre, Bombay-400 085, India.

loys [2-5] or on metal oxides [7] and a few studies on amorphous thin film alloys have also been reported [3]. Based on the kinetic analysis of the sputtering process in a binary alloy system, simple expressions have been proposed for determining the component sputter yield ratio [8,9]. The value of the component sputter yield ratio reflects the extent of preferential sputtering in a binary alloy target. Besides these simple kinetic analyses of the sputtering process, attempts have been made to understand the phenomenon of sputtering in terms of an interaction between the incident ions and the atoms of a solid target. In these investigations the observed compositional changes due to preferential sputtering have been correlated with the differences in the masses or in the chemical binding (i.e. the surface binding energies) of the alloy constituents [3,4]. Furthermore, the knowledge of the sputtering behaviour of a multicomponent

016%4332/91/$03.50 © 1991 - Elsevier Science Publishers B.V. All rights reserved

140 S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized Fe75B25

target is important in the quantification of data obtained by using AES or XPS in conjunction with ion sputtering [10].

The present study was motivated by the fact that it would be interesting to examine the sputter-induced surface compositional changes in a simple binary alloy in its amorphous and the corresponding crystalline states. The mass effects being the same, the differences in the chemical binding in the amorphous and the corresponding crystalline forms should be reflected in such an investigation. For this purpose, a melt-spun amorphous alloy of nominal composition FevsB25 was chosen. This alloy can be made to transform into a metastable single-phase Fe3B or a two-phase a-Fe + FezB structure on crystallization [11,12]. Therefore, a direct comparison of results in the amorphous and the crystalline forms of this alloy can be made without the complexities of sputtering in multiphase structures. Further, i ron-boron glasses constitute model systems from the point of view of determining structural information in metal-metalloid type glasses. These structural investigations have yielded information about the average atomic distances and the partial coordination numbers in these alloys [13-17]. Such information is important for correlating the results of preferential sputtering with chemical binding between alloy constituents as the latter depends on the distribution of nearest neighbours around an atom. It is thus expected that the extent of preferential sputtering in the amorphous and the corresponding crystalline forms may be different, in case the significant differences are introduced in the surface binding energies of the alloy constituents by the nearest neighbour environment of atoms in different structural states of the alloy.

No systematic investigation pertaining to the effect of preferential sputtering in a melt-spun amorphous alloy and its corresponding crystalline forms is available in the literature. Few previous investigations in amorphous systems have mostly been carried out on thin amorphous films [3] and some recent investigations using melt-spun amorphous alloys only mention the changes in the surface composition on sputter ion bombardment during the course of native oxide analyses [18,19]. The present study examines the effects of ion

bombardment on the surface composition in a melt-spun amorphous alloy Fe75B25 and its corresponding crystalline metastable Fe3B and the stable a-Fe + Fe2B structures. An attempt has been made to rationalize the observed changes in surface composition in terms of the expected differences in the surface binding energies in different structural forms of the alloy FevsB25. In addition to this, the native oxides formed on these different alloy forms of Fe75B25 have also been characterized by AES and XPS and an intercomparison of data is presented in the paper. Quite a few investigations pertaining to the analysis of native oxides, especially on F e -B based glasses, have been reported in the literature [18-22]. The characterization of surface oxide layers is important in understanding the role of the alloying elements in oxide film formation as it is known that additions like P, C, Si, B or Cr to Fe-based glasses dramatically affect the surface properties with regard to reactivity and corrosion resistance [23,24]. The interest in this area has grown further because of their potential for catalytic applications [25].

2. Experimental

Specimens, each of a suitable length ( - 15 mm), were cut from amorphous ribbon of FevsB25 (4 mm wide × 40 t~m thick) produced by the melt-spinning technique. The bulk composition of the ribbon was determined by using the techniques of X-ray fluorescence and atomic absorp- tion and showed 24.7 at% boron and 75.3 at% iron. The amorphous alloy FevsB25 exhibited a crystallization temperature (T×) of 740 K in a differential scanning calorimeter (DSC) at a a heating rate of 20 K/ra in . After mechanical polishing the shiny side (the air-side surface during melt-spinning) on a fine-grained silicon carbide paper, the specimens were electropolished in an electrolyte containing 90% ethanol and 10% per- chloric acid. The bath containing the electrolyte was maintained at 243 K and operated at a voltage of 20 V. After electropolishing, the specimens were cleaned under a jet of ethanol and dried in air. A few specimens after mechanical polishing were sealed in a capsule which was filled with

S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized Fe75B2s 141

spectroscopically pure argon gas at a pressure of about 220 Torr. These sealed specimens were heated in a furnace at a heating rate of 10 K / m i n and up to 793 K. The X-ray examination of these annealed specimens confirmed their structural state as the metastable crystalline Fe3B. A similar heat treatment has earlier been reported to lead to this phase transformation in this alloy [11,12]. Another batch of mechanically polished specimens was subjected to a heat treatment at 823 K in a high vacuum furnace for 2.5 h and the X-ray examination of these specimens after annealing showed the presence of a-Fe and Fe2B phases in them. These are the final stable phases obtained after the crystallization of amorphous Fe75B:5 alloy [11,12]. The crystallized specimens, consisting of metastable Fe3B and those consisting of a-Fe + FezB were electropolished as described earlier for amorphous specimens.

Quantitative AES analysis was carried out using a scanning Auger microprobe (PHI Model 600) equipped with a single-pass cylindrical mirror analyser. AES was performed at a primary beam voltage of 3 kV, electron beam current of 700 nA. The electron beam had a beam diameter of 3 #m and was incident at 30 o to the specimen surface normal. The beam was also rastered during the analysis over an area 50 /~mx 50 /~m on the specimen surface. The base pressure in the chamber was better than 1.0 × 10 - 9 Torr. The specimens were ion bombarded with 1 or 3 keV argon ions obtained from a differentially pumped ion gun. The ion beam was incident at an angle 56 ° to the surface normal and was rastered over an area 2 mm x 2 mm. The ion current densities for 1 and 3 keV Ar + beams were 14 and 40 /~A/cm 2, respectively. The native oxide films, formed prior to analysis on the specimen surfaces during an exposure of about 3 days in atmosphere after electropolishing, were also analysed. The depth profiles of native oxides were obtained using 3 keV argon ion having an ion current density of 2 /~A/cm 2 and the corresponding sputtering rate as obtained after sputtering a 100 nm thick Ta205 film was 0.46 nm/min . For quantitative analysis, both the low as well as the high energy Auger peaks of iron, namely Fe(47 eV, MVV) and Fe- (651 eV, LMM) along with the Auger peaks of

boron B(179 eV, KLL) were monitored. During depth profiling of native oxides, the Auger peaks of oxygen O(510 eV, KLL) and that of carbon C(271 eV, KLL) were also recorded. The Auger signals were measured in the normal mode and later differentiated on a computer interfaced with the system.

The XPS analysis of native oxide films was carried out on Perkin Elmer 5300 ESCA system using unmonochromatized X-rays from a Mg target (energy M g K a = 1253 eV). The system was equipped with a spherical capacitor analyser (SCA) and a position-sensitive detector (PSD). The base pressure in the chamber was better than 1.0 × 10-9 Torr. The size of the X-ray beam was about 2 mm × 10 mm and the spectra were recorded at a pass energy of 17.8 eV of the analyser. The subsurface layers were analysed after sputtering the surface with argon ions of 3 keV energy obtained from a differentially pumped ion gun. The ion beam was incident over an area of 1 cm × 1 cm on the specimen surface and yielded a sputtering rate of about 0.3 n m / m i n for a 100 nm thick Ta205 film. The XPS spectra were smoothened and deconvoluted using a computer interfaced with the system.

3. Data analysis and results

3.1. Quantification of A E S data

The measured Auger peak-to-peak heights were transformed to yield the surface composition by using the following expressions [26]:

g~ e IFe/IOe X~ I " / I o , (1)

F = [1 + rm(B)] [1 + rve(Fe)] Xm(B)~Fe(Fe)N°~

X ([1 + rm(Fe)] [1 + r , (B) ]

X ~k m (Fe) ~ , (B) N ° } - 1 , (2)

where I w and I B represent the Auger intensities from the binary alloy Fe75B25, I°~ and I ° represent the Auger intensities from the pure standards of iron and boron, respectively, N~ ° is the atomic density of i, Xj(i) and [1 + rj(i)] are the escape

]42 S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized FezsBes

Table 1 Values of various parameters used in eqs. (1) and (2)

Auger peaks X m(Fe) ~ Fe (Fe) R m(Fe) R re (Fe) F 1% used h m (B) h B (B) R m(B) R B(B) I ° F

Fe(47 eV) 0.35 0.36 1.79 1.87 1.12 1.91

B(179 eV) 0.59 0.48 1.66 1.28

Fe(651 eV) 1.12 1.14 1.46 1.51 1.10 1.11

B(179 eV) 0.59 0.48 1.66 1.28

Values of hi(j) are in nm and Ri(j) = [1 + ri(j) ].

depth and the backscattering factor for electrons from i in the matrix j. The subscript m refers to the matrix Fe-B alloy. The values of the escape depth ?~j(i) were calculated from the expression proposed by Seah and Dench [27] while those of the backscattering factor [1 + rj(i)] were evaluated using the expression of Ichimura et al. [28,29]. In evaluating these quantities in the alloy matrix (represented by m), the compositional averages of the atomic size and the atomic number were considered [30]. The atomic densities were taken from ref. [31]. The calculated values of various terms in eq. (2) for the case of the low energy Auger peak of iron Fe(47 eV) and the high energy Fe(651 eV)

Table 2 Values of the equilibrium surface composition ratio calculated from eq. (1) for 3 and 1 keV argon ion bombardments

Fe75B2 s a) 3 keV Ar + 1 keV Ar +

x~/ x ; / x~/ x~/ Xl%e c) X~ e d) Xl%e c) X~ e d)

Amorphous 0.439 0.330 0.430 0.327

Crystalline Fe3B 0.433 0.330 0 . 4 2 5 0.316

Crystalline et-Fe + Fe2B 0.374 0.309 0 . 3 5 1 0.290

Crystalline Fe2B b) 0.569 0 . 4 5 9 0 . 5 3 1 0.429

") The bulk composition ratio xb/Xbve = 0.328. b) X~/Xbe = 0.499. c) Using Fe(47 eV) and B(179 eV) Auger peaks. a) Using Fe(651 eV) and B(179 eV) Auger peaks.

Z 0

Z iii c.) z 0 ~J

C)

~E

0

100 a

i : ! i

50

0 0

.... = i I : , i . . . . . . . . . . I

20 Z.,O 60 80

SPUTTER TIME ( m i n i

1 0 0

z 0

(J Z 0 o

L)

:E 0

<C

50

0 0

; -i 20 Z.O 60 80

SPUTTER TIME ( m i n l

Fig. 1. Atomic concentration plots for the bombardment of amorphous Fe75B2s with Ar + ions of 3 and 1 keV energies: (a) using low energy Fe(47 eV) peak; (b) using high energy Fe(651 eV) peak.

S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized Fe75B25 143

Table 3 Experimental and calculated values of the sputter yield ratio YFe/YB using the values of the surface composit ion determined by Fe(47 eV) and B(179 eV) Auger peaks

Fe75B75 Experimental (from eq. (4)) Calculated for 3 keV Ar + bombardment (from eq. (15))

3 keV Ar + 1 keV Ar + Za ZF e [ref.] UB/UFe Yw/YB

Amorphous 1.34

Crystalline FeaB 1.32

Crystalline a-Fe + Fe2B 1.14

Crystalline FezB 1.14

6.89 1.31 12.91 [16] 0.873 1.18

9.0 1.30 13.7 [37] 0.965 1.30

1.07 - - -

8.0 1.06 15.0 [38] 0.860 1.05

peak are mentioned in table 1. Fig. 1 gives a typical plot of atomic concentration profiles from the amorphous FevsB25 specimen bombarded with the argon ions of 3 and 1 keV energies. The values of the surface composition ratio X~/X~e as calculated from eq. (1) for the bombardment of structurally different Fe75B25 alloys with 3 and 1 keV argon ions, are given in table 2. The surface composition ratio for the phase Fe2B was evaluated from the data obtained for a two-phase mixture of a-Fe and FezB in the following manner. The ratio of boron to the iron concentration fraction in Fe2B can be related to the corresponding ratio in the alloy consisting of the two-phase mixture of ct-Fe and Fe2B by the following expression:

= X~e .25 XFe Fe2B a-Fe + Fe2B

The above expression provided a simple approach to extract the data for Fe2B from those for a two-phase mixture of a-Fe and Fe2B. It should represent a reasonably good approximation in view of the fact that one of the phases in the two-phase alloy et-Fe + FezB is a pure elemental phase a-Fe whose concentration is assumed to remain con- stant during preferential sputtering. It would be worthwhile to mention here that the procedure employed for the quantification of Auger data yields an accuracy of 2% to 15% in composition measurement on non-sputtered surfaces [32]. The values of the surface composition calculated according to this procedure may be different from

those obtained by using experimentally known sensitivity factors derived from calibration standards of the same alloy. However, as the present measurements were made under identical condi- tions and were reproducible, it is possible to make an intercomparison of data obtained in different cases.

3.2. Component sputter yield ratio

The steady-state surface composition within the altered layer formed during ion bombardment is related to the bulk composition by [3,8]

S~e Y B S b e = YF~ X b ' (4)

where Y~ is the sputtering yield of the component i. The values of the sputter yield ratio Yve/Ys in different alloy forms of FevsB25 as evaluated from eq. (4) are mentioned in table 3.

3. 3. The chemical binding effects

For a comparison of the sputtering behaviour of the alloy Fe75Bz5 in different structural states, it would be interesting to examine the contribution to preferential sputtering arising due to the binding effects because the mass-dependent effects should be the same in all the cases. The binding effects appear due to the differences in the surface binding energies of the alloy constituents [3,4]. A thermodynamic formalism for

144 S.K. Sharma, S, Hofmann / A E S and X P S study of amorphous and crystallized Fe75B2s

evaluating the surface binding energy ratio in a binary alloy has earlier been proposed by Kelly [4]. This formalism is based on the pairwise nearest neighbour interactions between alloy constituents, but assumes that the bulk and the surface coordination numbers are the same for both types of atoms in the alloy. We would like to extend this approach for the case when the coordination of atoms around one type of atom is different from that around another type of atom in a binary alloy. This situation is most likely in amorphous alloys as has been revealed by structural investigations [13-16,33].

Consider a binary alloy A - B having surface atom fractions X~, X~ and bulk atom fractions X b, X b for the components A and B. Let Z A and Z B represent the total coordination numbers around an A atom and B atom, respectively. These can be further written in terms of the partial coordination number Z~j:

z~ = zA~ + ZAB, (5)

Z B = ZuB + ZUA, (6)

where Zij represents the number of " j " - t y p e atoms around an " / " - type atom. In terms of the pairwise nearest neighbour interaction formalism [34,35], the component surface binding energies can be expressed as:

2 Z A X , ~ U A A 2 "-~A'" B'-'AB' (7) UA ~ 1 t s _ _ l T P y s l T

1 t s 1 t s U B = -- ~ Z B XBUBB -- ~ Z B X A U B A , ( 8 )

where Z A and Z~ represent the respective coordination numbers on the surface in contrast to the similar quantities Z A and Z B in the bulk and U,j is the nearest neighbour bond strength for the bond i-j. Assuming a regular solution model, the quantities U, and U,j can be related to the heat of atomization A H, ~ for the pure element i and the heat of mixing A Hm for the alloy by the following expressions:

A H ~ = -- 1ZAUAA , ( 9 )

AH~] = ' (10) -- 7 Z B U B B ,

A H m = x b x b ( z ) [ U A B - - ½(UAA + UBB)] , ( 1 1 )

where ( Z ) denotes some average coordination number for the alloy [33]:

( Z ) = XbAZA + xbzB. (12)

We assume that the surface coordination numbers are one half of the respective bulk coordination numbers, i.e.,

Z A=½ZA, Z B = l Z B, etc. (13)

This is a reasonably good approximation in view of the fact that the atoms on the surface have the upper half of the plane of atoms missing. With this assumption and the approximation that the terms containing A H m do not make significant contributions, eqs. (7) and (8) can be simplified to yield the surface binding energy ratio UB/UA:

U s ZAZi~(1 + X~) + ZBX~AH~, Z B = Z~AH~(1 + X2) + ZAXaAH~ ZA" (14)

The above expression was used to evaluate the ratio (U~/Uve) using the experimentally determined values of the surface composition (see table 2) and the reported values of the heats of atomization AH, a [36] and the partial coordination numbers [13-16,37,38] in different alloy forms. The calculated values of UB/UA are mentioned in table 3.

3.4. Analysis of native oxide films

3.4.1. AES results Fig. 3 represents typical atomic concentration

depth profiles of the native oxide film on amorphous FevsBz5 for Fe(47 eV), B(651 eV), C(271 eV) and O(510 eV) Auger peaks. The depth profiles obtained from native oxides on crystalline Fe3B had similar features as found for the amorphous FevsB25. However, the native oxide depth profiles from the crystallized a-Fe + F%B alloy (shown in fig. 4) exhibited some differences which are discussed later in section 4.

3.4.2. XPS results The technique of XPS was used for identifying

the chemical states of elements present in native oxide films. The core level XPS peaks of iron, boron and oxygen were recorded during the native oxide film analysis on each of the specimens. Fig.


5, fig. 7 and fig. 8 depict the typical XPS peaks of Fe2p, O ls, and B l s, respectively, for the amorphous specimen. The corresponding XPS peaks from other specimens had similar features with some minor differences which are discussed in the next section.

4. Discussion

the surface contrary to the observations. There- fore, a mechanism which takes into account the binding effects must be invoked. The binding effects play a role in changing the surface composition through processes like cascade sputtering or bombardment-induced surface segregation [3,42, 45,49]. It has been shown that during cascade sputtering, the sputter yield ratio can be expressed as [31:

It can be observed from table 2 that the values of the surface composition calculated by using the low energy iron peak (47 eV) suggest an enrichment of boron on sputter bombardment of amorphous and crystallized specimens of FevsB2s with energetic argon ions. No systematic investigation pertaining to the sputtering behaviour of a melt-spun amorphous Fe -B and its corresponding crystallized forms has been reported in the literature. However, some similar observations in this regard have been made during the analysis of amorphous Fe -B based alloys [18,22]. A comparison with the values of the pure elemental sputtering yields [39-41] would also suggest the enrichment of boron on ion bombardment of a Fe -B alloy. But the predicted value of the sputter yield ratio YFe/YB on this basis is much larger, being about 4 to 5 [39-41], than that experimentally observed in the present study (1.06 to 1.34 based on the composition values obtained by using Fe- (47 eV) peaks, see table 3). The correction for the atomic density, as suggested by Ho et al. [9], to the experimental values of the sputter yield ratio re- suits in a still larger disagreement with the predicted values. Therefore, the pure elemental sputter yield values correctly predict the direction of enrichment, but not its magnitude.

Various mechanisms for understanding the preferential sputtering behaviour of alloy have been proposed and reviewed in the literature [3- 6,42-49]. These basically take into account the differences in masses or surface binding energies of the alloy components. A consideration of a mechanism on the basis of mass alone, e.g., as in recoil sputtering [44], does not explain the data obtained here as it gives a much lower value of the sputter yield ratio (YFJYB)=(MB/Mw) 1/4= 0.6633, thus suggesting the enrichment of iron on

YB + vx o UFe' (15)

where ~, = 4MveMB/(Mve + MB) 2 is the energy transfer factor and UB/UFe represents the surface binding energy ratio of boron to iron. The value of the surface binding energy ratio UB/UFe can be obtained from eq. (14) using the known values of the total partial coordination numbers [13- 16,37,38] and the heats of atomization [36]. As can be seen from table 3, there is a much better agreement between the calculated and the experimental values than that suggested by the predic- tions earlier. Cascade sputtering thus appears to be one of the probable mechanisms of sputtering in the present case. However, we have an interesting observation from table 2. It can be seen that the values of the equilibrium surface compositions calculated by using the low energy Auger peak of iron Fe(47 eV) are much different than those obtained from the high energy Fe(651 eV) peak and the latter in fact are close to the values in the bulk. A change in the energy of the argon ions indicates a similar trend and the surface compositions calculated using the high energy peak for 1 keV argon ion bombardment were even below the bulk values. It is noteworthy here that the high energy peak of iron probes more than three times the distance probed by the low energy peak (1.12 nm as against 0.35 nm, see table 1) [27]. It has been suggested that near room temperature the compositional changes in the alloy due to preferential sputtering extend to a depth approximately equal to the damage range of the bombarding ions [6,47,49]. The damage range in the present case was obtained by simulating the distribution of total target atom displacements using the TRIM programme [50] and was found to be - 2 . 5 nm


and - 5 . 0 nm for 1 and 3 keV argon ion bombardments, respectively. This calculation shows that the depth probed by the high energy Auger peak of iron (1.12 nm) lies well within the damage range of the bombarding ions. Therefore, the surface composition measured by the high energy Auger peak suggests a depletion of boron in the sub-surface region in comparison to the boron composition on the top surface region as probed by the low energy peak of iron. Such a situation is likely if surface segregation occurs during sputter ion bombardment. The problem of radiation-induced segregation during preferential sputtering at room temperature has been discussed by many authors [2-6,42-49]. Quite a few investigations have also been reported in the literature [51-54] which show the occurrence of Gibbsian- type segregation during sputter bombardment at room temperature. In general, near room temperature, where the point defect mobility is limited, preferential sputtering and displacement mixing are the main processes that govern the develop- ment of the alloy composition in the altered layer which extends to a depth approximately equal to the damage range. However, in addition to these, in alloys in which the point defects created during ion bombardment have sufficient mobility near room temperature a n d / o r Gibbsian adsorption of surface active elements is strong, substantial change in the near-surface composition can be produced as a result of bombardment-induced Gibbsian segregation. It can be seen from table 3 that U B < UF~, the segregation of weaker-bonded species boron will be favoured. Further, diffusion measurements carried out in boron-containing metallic glasses have suggested that the smaller boron atom diffuses by more than two orders of magnitude faster than the bigger iron atom [55,56]. In view of its much smaller size, an interstitial type of mechanism has been proposed for boron diffusion [57]. Therefore, the possibility of the occurrence of boron segregation to the surface during sputter-bombardment cannot be ruled out. In order to get some more information about the occurrence of such an effect from the present data, the following simple model calculation was attempted. It was assumed that the concentration distribution within the sputter damaged region

XO X2

x b

~_.- X Xl

a)

I I I I I I I !

zl 2zl Depth z (a.u.)

b)

~ x 1

e n . _ Xb Z X2

1 - x o

zl 2zi Depth z (o.u.)

Fig. 2. (a) Assumed concentration distribution of iron as described by eq. (16). The distribution for boron is obtained by

XB(z) =1 - XFe(Z) and is given in (b).

would look like as shown in fig. 2. Such a distribution can be expressed as:

xF (z)

= X ] f Fe, ( Xo- xb¢) e-(Z-z,'/L + xb~,

O ~ Z < Z l ,

Z>ZI,

(16)

where X ~ is the fraction of iron atoms in the first atom layer from the surface, )(be is the iron fraction in the bulk, z] is the thickness of the first layer (equal to the monolayer thickness), X 0 and L are constants. L represents the depth beyond which the sputter-induced effects are assumed to cease and can be approximated by the damage range of the bombarding ions in the alloy as mentioned earlier. The concentration distribution for boron can simply be written by using the fact that

XB(z ) = 1 - XFe(Z ) . (17)


The concen t ra t ion ra t io X~o/X~, de te rmined assuming a homogeneous concen t ra t ion wi th in the specimen, can be re la ted to the " t r u e " concent ra - t ion d i s t r ibu t ion b y the fol lowing express ion:

oo

f_ Xve(Z ) e -z/xF~ d z X~e = _ _ ~o ~ (18)

Xs XF° fo X s ( z ) e-~/xB d z

The rat io X~o/X~ was de t e rmined exper imen ta l ly us ing the low and the high energy Auge r peaks of i ron and is represen ted by (X~o/X~) e and (X~o/X~)n, respectively. Us ing the assumed concen t ra t ion d i s t r ibu t ion as given by eq. (16), eq.

(18) can be solved for (X~o/X~)e and (X~e//Xt~)h yie ld ing the fol lowing equat ions :

[(1 - e -d~) + Re(1 - - e - d 3 ) ] X~o

+ [r l e -a '+Rer3e - a 3 ] X 0

= [ ( r I - 1 ) e - d ~ + Re(r 3 - 1) e-d~]X~o + Re, (19)

[(1 - e - ~ ) + R , (a - e - ~ ) ] X k

+ [r= e - ~ + R~r~ e-"3] Xo

= [ ( r ~ - 1) e - ~ + R,(~3 - 1) e-~3] X~o + R , . (20)

where

R , = (X~e/X~)g, Rh = ( X ~ e / / l ~ ) h ,

d 1=zl /X1, d2 =z l /Xa , d3 = z J X 3 , L L L

r ~ = L + X ~ ' 1"2= L-4-)t 2 ' r 3 = LW~k 3 "

Here h~ ( i = 1 to 3) are the cor rec ted values of escape dep ths for Fe(47 eV), Fe(651 eV) and B- (179 eV) A u g e r peaks , respect ive ly (X~ = 0.64X, where X was ca lcu la ted as descr ibed ear l ier in the text), z~ was ca lcu la ted as 0.23 n m and L was a p p r o x i m a t e d by the d a m a g e range of ions ( L = 2.5 n m and 5.0 n m for 1 and 3 keV A r + ions, respect ively) . Eqs. (19) and (20) were s imul ta- neous ly solved for the unknowns X ~ and X0 using the k n o w n values of all the o ther pa rame te r s to y ie ld the concen t r a t i on f rac t ions in the first layer X ~ , X~ = (1 - X~¢) and in the second layer

x~o = ( Xo - x~o )e - ~ J L + X~o, x ~ = (1 - x~o). These values for the case of 3 keV A r + b o m b a r d - men t are m e n t i o n e d in tab le 4. The values calcu- la ted for 1 keV A r + b o m b a r d m e n t showed a more or less s imi lar t r end as for 3 keV and thus are not given here. I t is obse rved f rom this tab le tha t the first layer is s t rongly enr iched in b o r o n and its concen t r a t i on in the second layer is severely de- ple ted. In fact, the second layer appea r s to con ta in ha rd ly any b o r o n in the case of a m o r p h o u s and

Table 4 Calculated compositions of bombardment

the first and the second layers and the values of the sputtered yield ratio Yw/YB for 3 keV Ar + ion

Alloy 3 keV Ar + YFJ YB UB/UFe ( Y F e / Y B ) . . . . de

First layer Second layer (using X~e, X 1 (using X~¢, X~ (from eq. (15) in eq. (4)) in eq. (14)) and UB/UF¢ from

previous column)

Yro/xB ~)

Fe75B25 XIe = 0.42 X~c ~ 0.99 4.21 0.82 1.11 1.28 (amorphous) X~ = 0.58 X~ = 0.01

Fe3B X~e = 0.42 Xge = 0.99 4.21 0.93 1.26 1.28 (crystalline) X~ = 0.58 X~ = 0.01

a-Fe + FezB X~e = 0.54 X~e = 0.92 2.60 - 1.13 (crystalline) X~ = 0.46 Xa 2 = 0.08

Fe2B X#e = 0.46 X~e = 0.84 2.35 0.83 1.01 1.08 (crystalline) X~ = 0.54 Xs 2 = 0.16

., rFo = [ 9 - z) x~ + +x~ ] x~o Ys [ (1 - 3' ) X1Fe + "tX2¢ ] -X-B-s ' where the sputtered atom fraction ,/ from the second layer is assumed to be 0.5.

148 S.K. Sharrna, S. Hofmann / AES and XPS study of amorphous and crystallized Fe75B25

crystallized Fe3B. Therefore, this simple model calculation is suggestive of a strong boron segregation to the top surface layer and its corresponding depletion in the second and the subsequent layers. However, we would like to mention that this inference needs a conclusive experimental proof and further investigations using especially the technique of ion-scattering spectrometry (ISS) at low as at high temperatures must provide useful information in this regard. Based on the results of the present study, we would simply mention that cascade sputtering and f o r surface segregation appear to be the most probable processes in altering the surface composition on sputter-ion bombardment of amorphous and crystallized Fe75B25 alloys.

It is noteworthy from table 3 that the calculated values of the sputter yield ratio (Yve/YB) are not much different from the measured values. In this context, we would like to mention that this striking closeness between the measured values and those calculated from eq. (15) based on the cascade theory could be coincidental in view of the following. If the bombardment-induced segregation is the dominant process as suggested by the model calculation in the preceding paragraph, then the calculated values of the first atom layer would yield much higher values of the sputter yield ratio YFe/YB according to eq. (3) (e.g. Yve/Yn = 4.21 for amorphous FevsB25, see table 4). These values are much larger than those obtained from eq. (15) for cascade sputtering (e.g. Yw/YB = 1.11 for amorphous FevsB25 see table 4). This is not surprising in view of the fact that the available data on sputtering in crystalline alloys do also show a similar trend [3]. In fact, both the cascade sputtering and the surface segregation strongly correlate with the bonding, eq. (15) based on cascade theory correctly predicts the enriched species, but is unable to explain the magnitude of enrichment, if the Gibbsian-segregation is the dominant process during sputtering [3,45]. More- over, if a significant fraction of sputtered atoms arise from the second layer, in such a case the true value of the sputtered yield ratio would be determined by the compositions of the first as well as the second layer [6,49]. Assuming that the sputtered atom flux consists of 50% of atoms from

the second layer [6], the calculated compositions of the first and the second layers give values of the sputter yield ratio Yve/YB which do not differ much from the corresponding experimentally measured values (e.g., YFJYB = 1.28 from table 4 as against 1.34 from table 3 for amorphous Fe75B25 ). In this manner, it is possible to reconcile the differences between the experimentally observed values of YFJYB and those obtained from the calculated values of the composition of the first and the second layers. However, in the absence of measurements representing the true compositions of the first a n d / o r the second layers, the present measurements yield only the average values of YFJ YB over a much larger depth (e.g., about three layers using the surface composition values measured by low energy Fe(47 eV) and B(179 eV) Auger peaks.

An intercomparison of data mentioned in tables 2 and 3 indicates that the values of the equilibrium surface composition and the sputter yield ratio are quite similar in the amorphous and the corresponding crystallized Fe3B alloys. On the other hand, these values for the two-phase alloy a-Fe + FezB or those derived for FezB (see section 3.1) are much different from the values for the amorphous or the crystallized Fe3B alloys. It would be worthwhile to point out here that the mass effects should make a similar contribution to preferential sputtering in these alloys, and the observed differences in their preferential sputtering behaviour can be attributed to the binding effects. We have earlier mentioned that the binding effects make a significant contribution to the preferential sputtering of these alloys through cascade sputtering a n d / o r surface segregation. The binding effects, in turn, would be governed by the atomic structure of the alloy. Therefore, the observation of similar values of the equilibrium surface composition and the sputter yield ratio in the amorphous Fe75B25 and the corresponding crystallized metastable Fe3B would suggest that the local environment of an atom in the amorphous and the crystallized Fe3B are very similar. Similar conclusions have been drawn regarding the struc~ ture of amorphous Fev3B2s on the basis of investigations by M~Sssbauer spectroscopy [17,58,59] and by XPS investigations of the valence band of the


amorphous and the crystallized Fe3B [60]. More- over, the displacement threshold energy of iron atoms in amorphous and crystalline Fe3B alloys has been found to be the same [61]. On the other hand, the different values of these quantities (X~/X~ and YFe/YB) in the case of the two-phase a-Fe + Fe2B alloy or for FeEB (see tables 2 and 3) can be understood in view of the different structure of this alloy than that for FeaB [38,58,59].

Another noteworthy observation which can be made from table 3 is that the deviation between the calculated and the experimental values of YFe/YB is much smaller in the case of crystalline FeaB and the two-phase a-Fe + Fe2B alloys than in the case of the amorphous Fe75Bzs. We have earlier noted (see table 3) that the experimental values of YFe/YB are quite similar for the amorphous and crystalline FeaB. It is an interesting observation and suggests that the set of coordination numbers [13-16] as used in evaluating UB/UF~ according to eq. (14) for the amorphous FevsB25 may not be a true representative of the structure of the amorphous alloy. These coordination numbers have been obtained on the basis of a dense random packing of hard spheres (DRPHS) model for the amorphous alloy [13-16,62]. DRPHS has been considered a successful model for metal-metalloid type amorphous alloys as it explains many of their structure-sensitive properties [15,16]. Some investigations, especially those using the M~Sssbauer spectroscopy have suggested a quasi-crystalline model for the structure of Fe-B glasses on the basis of similarities found for the amorphous and the crystalline Fe-B alloys [58,59]. Our results on the preferential sputtering of amorphous and crystallized FeaB also corroborate this suggestion for the structure of amorphous Fe75B25. A quasi-crystalline model for the amorphous Fe75B25 would suggest a similar local order and hence the coordination of atoms as in the case of crystalline Fe3B [59]. Assuming a quasi-crystalline model for the amorphous alloy, the contribution to preferential sputtering due to binding effects can be expected to remain the same in both the amorphous and the crystallized Fe3B alloys. This is in conformity with the experimental observations which suggest similar values of the equilibrium surface composition and the

Ill: 60 . . . . .

' ° I ....... ......... i .......... ............... ! :, .........

t,, . . . . . . . . . . . . o ,,o . . . . . . . . . . . . . .

o 0 4 8 12 16 20

SPUTTER TIME (min

Fig. 3. A t o m i c concen t r a t i on d e p t h prof i les o b t a i n e d f rom the

na t ive ox ide f i lm on the a m o r p h o u s Fe75B25.

sputter yield ratio in both the amorphous and the crystalline Fe3B (see tables 2 and 3). Therefore, it is clear from the above discussion that the preferential sputtering behaviour of amorphous and crystalline Fe75B25 alloys can be understood in terms of the atomic structures of these alloys.

We shall now briefly and qualitatively discuss our results on the native oxide analysis by AES and XPS. The AES depth profiles of native oxide films formed on the amorphous and crystallized Fe3B were quite similar (fig. 3) while those for the two-phase t~-Fe + Fe2B alloy possessed some con- trasting features (fig. 4). An observation from figs.

z: Z o I - - < ¢r" I - -

Z LU t . ) Z o c_)

o

3r o i - - <

60

~0

20

j - - - -

. .----r.2 . r e _.

---"~"V J , _ - - T

0

", C

0 6 12 18 24 30

SPUTTER TiME (m in )

Fig. 4. A t o m i c concen t r a t i on p lo ts of na t ive ox ide f i lm on the two-phase a - F e + Fe2B alloy.

150 S.K. Sharma, S. Hofrnann / AES and XPS study of amorphous and crystallized FezsB25

3 and 4 would suggest that the oxide films were enriched in boron, the maximum enrichment being near the oxide /a l loy interface. The enrichment of boron was much more in the case of amorphous and crystalline Fe3B alloys than in the two-phase alloy. These observations are consistent with those reported in the literature [19-22]. The enrichment of boron near the oxide alloy interface suggests that the initial oxidation begins with the formation of a boron-rich layer. It is quite likely in view of a smaller formation enthalpy for boron (AH298K = -204 .07 kca l /mol O z for B203) than for iron ( A H = -126.40 kca l /mol 02 for FeO). The formation of this boron-rich layer in the interracial region can retard further oxidation of iron if it is uniform and continuous [64,65]. This appears to be true for the single-phase amorphous and the crystalline Fe3B alloys because the thicknesses of the oxide films formed on these alloys are similar and small (3.6 and 3.3 nm) as com- pared to that (9.7 nm) on the other alloy consisting of two phases (et-Fe + Fe2B ). It should be noted here that the values of the oxide film thickness are only approximate ones as these have been obtained on the basis of the sputtering rate for TazO 5. However, a qualitative comparison using these values can still be made. The larger thickness on the two-phase alloy indicates faster oxidation of this alloy in comparison to the amorphous and the single-phase Fe3B alloys. Wei and Cantor [64,65] have suggested, on the basis of their studies on the oxidation of Fe-based amorphous and their crystalline counterparts, that the oxidation behaviour of these alloys has a strong dependence on the alloy microstructure. A continuous and more uniform initial oxide film is more likely to form on the single-phase amorphous and crystalline Fe3B alloys while the formation of such a layer is prevented on the other alloy consisting of two phases a-Fe and Fe2B, thus explaining its faster oxidation. In fact, the native oxide depth profiles for the amorphous FevsB25 as shown in fig. 3 suggest a strong segregation of boron to the interfacial region with its corresponding depletion below the interfacial region during initial oxidation of this alloy. On the contrary, the depth profiles for a-Fe + Fe2B alloy as shown in fig. 4 would suggest a continuous diffusion of iron and boron

\ a)

c)

730 720 718 700

II]I~ING [141~, eq

Fig. 5. Fe2p XPS peaks from the native oxide film on amorphous Fe75B2s: (a) as-received; (b) sputtered for 1 min; (c) sputtered for 4 min; (d) sputtered for 12 min. Sputter ion

(Ar ÷ ) energy = 3 keV.

through the initial oxide film resulting in a much faster oxidation in this alloy.

We shall now discuss the XPS results. Fig. 5 shows a typical Fe2p XPS peak from the native oxide on amorphous Fe75B25 alloy before and after sequential sputtering with argon ions. The peak before sputtering appeared at 710.5 eV with a satellite peak at 719.8 eV while after sputtering the peak shifted to 710.0 eV, the satellite peak at

S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized FezsBes 151

719.8 eV disappeared and a new satellite peak at 715.8 eV appeared. A shoulder at 706.6 eV representing elemental iron was also present in all the spectra and became prominent after sputtering. A deconvolution procedure carried out on this peak (fig. 6) suggested the presence of peaks at 711.1, 709.2, and 706.6 eV representing the contributions from Fe 3+, Fe E+ and Fe ° species, respectively [66,67]. The film on the as-received surface consisted of more of Fe 3+ along with Fe 2+ and little Fe °. The peaks due to Fe 2+ and Fe ° grew in intensity after sputtering while that due to Fe 3+ became less intense. It is known that a partial conversion of Fe 3+ to lower valent species might occur due to sputtering effects [68]. The spectral features of the Fe 2p peak from all the specimens had somewhat similar characteristics with the peaks due to various species occurring within + 0.2 eV at positions mentioned above in the deconvoluted spectra. The contributions arising due to various species showed some variations in spectra from different specimens: the peak due to Fe 3+ was much bigger than that due to Fe 2÷ before sputtering in the two-phase a-Fe + Fe2B alloy while those from the amorphous and Fe3B specimens were of comparable intensities. Also the peak due to Fe ° was more intense in the

amorphous and Fe2B alloys than in the two-phase alloy. It is possible to understand this variation in contributions form various species in terms of the variation in their oxide film thicknesses which was maximum (9.7 nm) for et-Fe + Fe2B alloy and had comparable values of 3.6 and 3.3 nm in FeaB and the amorphous alloys, respectively. This analysis showed that the native oxide films contained Fe 3+, Fe 2+ and Fe ° (a small peak in the as-received spectra) in varying proportions and Fe 3+ species were predominant in the two-phase alloy a-Fe +

Fe2B. The XPS spectrum of boron (B ls peak, see fig.

7) showed a much bigger peak at 191.9 eV and a small peak at 187.6 eV. The former corresponds to boron in oxide form (possibly as B 3+ species) while the latter arises due to elemental boron [69]. The peak due to oxidic boron increased in irrten- sity in the sputtered spectra which is consistent with the results of AES depth profiles showing enrichment of boron in the interfacial region (see fig. 3). Further, the intensity ratio of B ls to Fe2p peaks was much larger for the amorphous and the crystalline Fe3B alloys than for a-Fe + Fe2B alloy. This observation is in line with the inference drawn from the native oxide depth profiles for these alloys (figs. 3 and 4). After prolonged sputtering,

10 I I I I I I I I ~ ^ I I I I I

9 Fe2p

7

4 /

3 A

718 716 71'1 712 710 70~ 706 BINDING ENERGY, eV

Fig. 6. Deconvoluted spectra in the Fe 2p spectrum corresponding to that shown in fig. 5c. The peaks due to various species are indicated in the figure.


Bls

S

v

w

c)

r/ /,

Fig. 7. B l s X P S peaks from the native oxide film on amorphous Fe75B25: (a) as-received; (b) sputtered for 1 rain; (c) sputtered for 4 rnin; (d) sputtered for 12 rain. Sputter ion ( A r + ) energy

= 3 keV.

the peak due to oxidic boron disappeared and only a single peak at 187.6 eV representing elemental boron remained. While studying the oxidation of Fe -B glasses, Nagarkar et al. [21] have observed that the boron peak occurs at 191.5 eV. These values are very close to our values, but are lower than the reported value at 193.1 eV for B203 [69]. In this context, Nagarkar et al. [21] have suggested that the boron peak at this lower posi-

tion corresponds to iron-boron oxides or to BxOy suboxides. On the other hand, Myhra et al. [22] have mentioned that the boron oxide peak could appear at lower positions in thin amorphous films.

The oxygen O ls XPS peak (shown in fig. 8) in all the cases had two principal components, one occurring at 531.2 + 0.1 eV and the other at 530.0 _ 0.2 eV. In addition to this, a broad shoulder at the high binding energy side was also present which could be accounted for in the deconvolution by fitting a peak at 532.6 _ 0.2 eV (see fig. 9). The spectra before sputtering showed that the peaks at 531.2 + 0.1 eV and at 530.0 + 0.2 eV were of comparable intensities in the amorphous specimen while the former was much bigger than the latter in the crystallized specimens. On sputtering these

f J

J /

/

.x~ '

~ L

/ //

Ols

L\

\,

'\ b )

/

/ /

540 536 532

BINDING ENEI~Y, eV

, c)

528

Fig. 8. O ls XPS peaks from the native oxide film on amorphous Fe75B25: (a) as-received; (b) sputtered for 1 min; (c) sputtered

for 4 min.

S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized FezsB25 153

v z

10

9

8

?

6

5

4

3

2

t

0

I I I I I I I I ~ I I / ~ l I I I ~

536 535 534 533 532 531 530 529 528 BINBIN6 EER6Y, eq

Fig. 9. Deconvoluted spectra in the O l s spect rum corresponding to that shown in fig. 8a. The peaks due to various species are indicated in the figure.

two peaks had comparable intensities in the two- phase alloy while in the amorphous and the Fe3B alloys, the component at 531.2 + 0.1 eV was much more intense than the other at 530.0 + 0.2 eV. The peak at 531.2_+ 0.1 eV corresponds to hydroxyl ion ( O H - ) bonding while that at 530.0_ 0.2 eV represents metal-oxide ( 0 2- ) bonds [66,67]. It has been reported that the O ls peak for B203 appears at 533.0 eV [69]. The shoulder at the high binding energy side corresponding to a peak at 532.6 eV can be attributed to boron-oxygen bonds.

Therefore, a combined analysis of Fe2p, O ls and B ls XPS peaks would suggest that the native oxide films on all the specimens consisted of hy- droxides and oxides of iron (possibly Fe(OH)3, Fe(OH)2, Fe304), in addition to boron oxide (as B203 or as suboxides) and some iron in unoxidized form also. The as-received surface also contained Fe(III) oxide (possibly Fe203) as indicated by a small satellite peak at 719.8 eV. The formation of FeOOH on the surface is not suggested because the component representing the O H - peak (531.2 eV) in the O ls doublet structure is known to show' dramatic reduction on ion bombardment [66] contrary to the observations in the present study. Our results are different from those of Myhra et al. [22] who carried out XPS investigations of the native oxide films formed on amorphous F e - B

alloys. They suggested that the native oxide film mainly consisted of Fe 3+ species (as Fe203). The differences between their results and those in the present study become obvious because Myhra et al. [22] analysed the native oxides on the as- quenched specimens without any pre-polishing while the specimens used in our study were electropolished prior to the formation of native oxide films. Formation of Fe203 on the surface is most likely to occur during glass formation itself as the glass is formed at a high but finite cooling rate and the formation of F e 2 0 3 may be favoured if the glass remains at a higher temperature even for a very short fraction of time during melt spinning in the presence of oxygen-containing atmosphere.

4. Conclusions

The present investigation, dealing with the Study of the preferential sputtering of the amorphous and crystallized FeTsB25 (the metastable Fe3B and the stable two-phase mixture of a-Fe + Fe2B) alloys and with the characterization of the native oxide films on these' alloy forms, revealed the following: (i) The preferential sputtering resulted in the formation of a boron-rich sputtered altered layer. A

154 S.K. Sharrna, S. Hofmann / AES and XPS study of amorphous and crystallized FezsB25

mode l ca lcula t ion based on the surface compos i - t ion measured by using the low Fe(47 eV) and the high energy Fe(651 eV) peaks of i ron suggested a s t rong b o r o n enr ichment in the top surface layer wi th its co r re spond ing dep le t ion in the subsurface layers. I t is suggested that the poss ib le spu t te r ing mechan i sms would include cascade spu t te r ing a n d / o r surface segregat ion in which the b ind ing effects p lay a s ignif icant role. (ii) The a m o r p h o u s Fe75B25 and the crys ta l l ized Fe3B d i sp lay a s imi lar sput te r ing behav iou r in con t ras t to that possessed by the two-phase ct-Fe + Fe2B alloy. I t is poss ib le to ra t ional ize these dif ferences in their sput te r ing behav iou r in terms of the differences in their s t ruc ture and are sugges- t ive of a s imilar local o rder in the s t ructures of the a m o r p h o u s Fe75B25 and the co r r e spond ing crysta l - l ized Fe3B alloys. Based on the desc r ip t ion of the sput te r ing behav iou r of the a m o r p h o u s Fe75B25 p rov ided by the avai lab le s t ruc tura l in fo rmat ion , a quas i -crys ta l l ine mode l ra ther than a dense ran- d o m pack ing mode l appea r s to be a be t te r repre- senta t ive of the s t ructure of the a m o r p h o u s Fe75B25 alloy. (iii) The nat ive ox ide f i lms on these s t ruc tura l ly d i f ferent al loys forms were enr iched in b o r o n near the o x i d e / a l l o y in ter rac ia l region and cons is ted of oxides and hydrox ides of bo th i ron and boron . The ox ida t ion was rela t ively much faster in the t w o - p h a s e a - F e + Fe2B a l loy t h a n in the a m o r p h o u s and the s ingle-phase crys ta l l ized Fe3B alloys. This was a t t r i bu ted to the dif ferences in their micros t ruc tures .

Acknowledgements

Thanks are due to Mr. B. Siegle for the A E S analysis , to Ms. K. K r a u s for XPS analysis and to Mr. R.V. R a m a n a n of Al l i ed C o r p o r a t i o n for p ro- v id ing the a m o r p h o u s a l loy used in this invest iga- t ion. One of the au thors (S.K.S.) wou ld l ike to express his thanks to the Max-P lanck-Gese l l s cha f t for the award of a research fel lowship. Thanks are due to Dr. V. N a u n d o r f for p rov id ing access to the T R I M p r o g r a m m e and for m a n y useful discus- sions.

References

[1] S. Hofmann, Surf. Interf. Anal. 2 (1980) 148. [2] G. Betz, Surf. Sci. 92 (1980) 283. [3] R. Kelly, Surf. Sci. 100 (1980) 85. [4] R. Kelly, Nucl. Instr. Meth. 149 (1978) 553. [5] R. Shimizu, Nucl. Instr. Meth. B 18 (1987) 486. [6] N.Q. Lam, Surf. Interf. Anal. 12 (1988) 65. [7] J.B. Malherbe, S. Hofmann and J.M. Sanz, Appl. Surf.

Sci. 27 (1986) 87. [8] H. Shimizu, M. Ono and K. Nakayama, Surf. Sci. 36

(1973) 817. [9] P.S. Ho, J.E. Lewis, H.S. Wildman and J.K. Wonward,

Surf. Sci. 57 (1976) 393. [10] S. Hofmann, Depth profiling in AES and XPS, in: Auger

and X-Ray Photoelectron Spectroscopy, Vol. 1 of Practi- cal Surface Analysis, 2nd ed., Eels. D. Briggs and M.P. Seah (Wiley, Chichester, 1990) ch. 4, pp. 143-199.

[11] U. K/3ster and U. Herold, in: Glassy Metals I, Vol. 46 of Topics in Applied Physics, Eds. H.-J. Giintherodt and H. Beck (Springer, Berlin, 1981) p. 225.

[12] E. Kneller, Y. Khan, V. Geiss and J. Vosk~tmper, Z. MetaUkd. 75 (1984) 698.

[13] T. Fujiwara, H.S. Chen and W. Waseda, in: Proc. 4th Int. Conf. on Rapidly Quenched Metals, Vol. 1, Eds. T. Masumoto and K. Suzuki (Japan Institute of Metals, Sendai, 1982) p. 287.

[14] Y. Waseda and H.S. Chen, Phys. Status Solidi (a) 49 (1978) 1.

[15] Y. Waseda, Prog. Mater. Sci. 26 (1981) 1. [16] D.S. Boudreaux, Phys. Rev. B 18 (1978) 4039. [17] T. Kemeny, I. Vincze, B. Fogarassy and S. Arajs, Phys.

Rev. B 20 (1979) 476. [18] D.R. Huntley, S.H. Overbury, D.M. Zehner, J.D. Budai

and W.E. Brower, Jr., Appl. Surf. Sci. 27 (1986) 180. [19] P. Kulkarni, S. Sathe, M.G. Thube, A.S. Nigavekar and

S.K. Kulkarni, Appl. Surf. Sci. 37 (1989) 419, and references cited therein.

[20] S. Myhra and J.C. Riviere, J. Non-Cryst. Solids 99 (1988) 244.

[21] P.V. Nagarkar, S.K. Kulkarni and E. Umbach, Appl. Surf. Sci. 29 (1987) 194, and references cited therein.

[22] S. Myhra, J.C. Riviere and A.M. Stewart, J. Non-Cryst. Solids 99 (1988) 255.

[23] K. Hashimoto, in: Amorphous Metallic Alloys, Ed. F.E. Luborsky (Butterworths, London, 1983).

[24] K. Hashimoto, K. Asami, M. Naka and T. Masumoto, Corros. Sci. 19 (1979) 857.

[25] G.K. Kisfaludi, K. Lazar, Z. Sehay and L. Guczi, Appl. Surf. Sci. 24 (1985) 225.

[26] M.P. Seah, Vacuum 36 (1986) 399. [27] M.P. Seah and W.A. Dench, Surf. Interf. Anal. 1 (1979) 2. [28] S. Ichimura, R. Shimizu and J.P. Langeron, Surf. Sci. 124

(1983) IA9. [29] R. Shimizu, Jpn. J. Appl. Phys. 22 (1983) 1631. [30] C.G.H. Walker, D.C. Peacock, M. Prutton and M.M. E1

Gomati, Surf. Interf. Anal. 11 (1988) 266.

S.K. Sharma, S. Hofmann / AES and XPS study of amorphous and crystallized FezsB25 155

[31] C. Kittel, Introduction to Solid State Physics, 6th ed. (Wiley, New York, 1986) p. 24.

[32] H.J. Mathieu and D. Landolt, Surf. Interf. Anal. 3 (1981) 153.

[33] G.S. Cargill III, J. Non-Cryst. Solids 43 (1981) 91. [34] R.A. Swalin, Thermodynamics of Solids (Wiley, New

York, 1972). [35] F.L. Williams and D. Nason, Surf. Sci. 45 (1974) 377. [36] R. Hultgren, P.D. Desai, D.T. Hawkins, M. Gleister, K.K.

Kelly and D.D. Wagman, Selected Values of the Thermo- dynamic Properties of the Elements (American Society for Metals, Metals Park, OH, 1973).

[37] P. Lamparter, W. Sperl, E. Nold, G.-R. Harbach and S. Steeb, in ref. [13], p. 343.

[38] E.E. Having, H. Damsma and P. Hokkeling, J. Less-Com- mon Metals 27 (1972) 169.

[39] M.P. Seah, Thin Solid Films 81 (1981) 279. [40] N. Matsunami et al., Energy Dependence of Sputtering

Yields of Monoatomic Solids, Report No. IPPJ-AM-14, Institute of Plasma Physics, Nagoya University, Chkusaku, Nagoya, Japan, June 1980.

[41] H.H. Andersen and H.L. Bay, in: Sputtering by Particle Bombardment I, Vol. 47 of Topics in Applied Physics, Ed. R. Behrisch (Springer, Berlin, 1981) p. 145.

[42] G. Betz and G.K. Wehner, in: Sputtering by Particle Bombardment II, Vol. 52 of Topics in Applied Physics, Ed. R. Behrisch (Springer, Berlin, 1983) p. 11.

[43] P. Sigmund, in ref. [41], p. 9. [44] R. Kelly and D.E. Harrison, Mater. Sci. Eng. 69 (1985)

449. [45] R. Kelly, Surf. Interf. Anal. 7 (1985) 1. [46] D.G. Swartzfager, S.B. Ziemecki and M.J. Kalley, J. Vac.

Sci. Technol. 19 (1981) 185. [47] J. Kirschner, Nucl. Instr. Meth. B 7 /8 (1985) 742. [48] P.H. Holloway and R.S. Bhattacharya, J. Vac. Sci. Tech-

nol. 20 (1982) 444. [49] N.Q. Lam and H. Wiedersich, Nucl. Instr. Meth. B 18

(1987) 471.

[50] J.P. Biersack and L.G. Haggmark, Nucl. Instr. Meth. 174 (1980) 257.

[51] T. Koshikawa and K. Goto, Nucl. Instr. Meth. B 18 (1987) 504.

[52] G. Hetzendorf and P. Varga, Nucl. Instr. Meth. B 18 (1987) 501.

[53] J. Du Plessis, G.N. Van Wyk and E. Tagiauer, Surf. Sci. 220 (1989) 381.

[54] R.-S. Li, C.-F. Li and W.-L. Zheng, Appl. Phys. A 50 (1990) 169.

[55] R.W. Cahn, J.E. Evetts, J. Patterson, R.E. Somekh and C.K. Jackson, J. Mater. Sci. 15 (1980) 702.

[56] K. Stobiecki, W. Palmer, W. Reill, K. RSll and H. Hof- mann, J. Non-Cryst. Solids 88 (1986) 209.

[57] B. Cantor and R.W. Cahn, in: Amorphous Metallic Al- loys, Ed. F.E. Luborsky (Butterworths, London, 1983) p. 487.

[58] I. Vincze, T. Kemeny and S. Arajs, Phys. Rev. B 21 (1980) 937.

[59] I. Vincze, D.S. Boudreaux and M. Tegze, Phys. Rev. B 19 (1979) 4896.

[60] M. Matsuura, T. Nomoto, F. Itoh and K. Suzuki, Solid State Commun. 33 (1980) 895.

[61] A. Audouard, J. Balogh, J. Dural and J.C. Jousset, Rad. Eft. 62 (1982) 161.

[62] D.E. Polk, Acta Met. 20 (1972) 485. [63] M.G. Thube, C.V. Dharmadhikari, G.Y. Dixit, S.K.

Kulkarni and A.S. Nigavekar, Surf. Sci. 182 (1987) 439. [64] G. Wei and B. Cantor, Acta Met. 36 (1988)1 67. [65] G. Wei and B. Cantor, Acta Met. 36 (1988) 2293. [66] C.R. Brundle, T.J. Chuang and K. Wandelt, Surf. Sci. 68

(1977) 459. [67] K. Wandelt, Surf. Sci. Rep. 2 (1982) 75. [68] D.F. Mitchell, G.I. Sproule and M.J. Graham, J. Vac. Sci.

Technol. 18 (1981) 690. [69] D.J. Joyner and D.M. Hercules, J. Chem. Phys. 72 (1980)

1095.

study of amorphous and crystallised fe75b25 by auger electron spectroscopy and x-ray photoelectron...

Documents